DEVELOPMENTS IN lHE SCIENCE

AND TECHNOLOGY OF COMPOSrTE

MATERIALS

E~M EUROPEAN ASSOCIATJON FOR COMPOSITE MATERIALS

DEVELOPMENTS IN THE SCIENCE

AND TECHNOLOGY OF COMPOSITE

MATERIALS

THIRD EUROPEAN CONFERENCE ON COMPOSITE MATERIALS

20.23 MARCH 1989 BORDEAUX-FRANCE

EDnTJRS: AR. BUNSELL, P. LAMICQ, A MASSIAH

ELSEVIER APPUED SCIENCE: LONDON AND NEW-YORK

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SCIENTIFIC COMMITTEE ECCM-3

President: Mr P. LAMICQ Societe Europeenne de Propulsion (SEP)

France

· SELECTION PAPERS COMMITTEE:

FRANCE A.R. BUNSELL

F.X.

Ecole Nationale Supcrieure des Mines de Paris

ITALY G.DiDRUSCO A.SAVAOORI

de CHARENIENA Y P.S.A. R. NASLAIN Laboratoire des Composites

Thermostructuraux

BELGIUM N.SPRECHER 1. VERPOEST

DENMARK H.LILHOLT

GREAT ·BRIT AIN M.G. BADER

SWEDEN Owens·Coming Fiberglas Europe Th. JOHANNESSON Kath Universiteit Leuven

Riso National Laboratory SWITZERLAND R.PINZELLI

THE NETHERLANDS

Montedison Enichem

Institute of Technology of Linkoping

Du Pont

de Nemours

J.H. GREENWOOD B. HARRIS

University of Surrey Era Technology University of Bath

I.A.N. SCOTT Shell

F.L. MATTHEWS

WEST·GERMANY

Imperial College of Science and Tcchnology

K. FRIEDRICH Technische Universitlit Hamburg-Harburg

G. GRUNINGER DFVLR H.KELLERER MBB

• ORGANIZING COMMITTEE: EACM

A.R. BUNSELL A. MASSIAH J.L. ZULIAN D. DOUMEINGTS

• CONGRESS SECRETARIAT:

H.BENEDIC C.MADUR

Laboratorium

PREFACE

Dr A.R. BUNSELL President de

l'Association Europeenne des Materiaux Composites

Apres Ie succes des deux precedentes editions, Ie troisieme Congres Europeen sur les Materiaux Composites, ECCM-3, s'annonce de tres haut niveau.

La premiere reunion s'est tenue a BORDEAUX en septembre1985 tandis que la seconde, jumelee avec Ie sixieme Congres International sur les Materiaux Composites, ICCM-6 eut lieu a LONDRES en juillet 1987. Ces deux reunions ont clairement montre I'importance de la recherche sur les materiaux composites en Europe. Elles ont pu rassembler les chercheurs venus de toute l'Europe et du monde entier.

Ce troisieme congres nous ramene a BORDEAUX et souligne I'interet porte aux materiaux composites a BORDEAUX et en Aquitaine, haut lieu d'application des technologies de pointe. La creation a BORDEAUX de l'Association Europeenne des Materiaux Composites, A.E.M.C., a ete Ie fait a la fois d'une volonte politique et d'une prise de conscience locales ainsi que de la presence d'un tissu industriel favorable au developpement de celie activite.

Le travail assidu du comite scientifique, compose de specialistes europeens, grace a qui les sujets traites et les articles sont varies et de haut niveau, nous assure de la qualite du compte rendu de ECCM-3 qui comptera desormais parmi les ouvrages de reference. Plus notable encore, est Ie fait que ECCM-3 contribue a la construction d'une grande communaute internationale des composites, composee de chercheurs venus du monde entier.

D'ores et deja, je vous donne rendez-vous pour Ie 4e Congres Europeen des Materiaux Composites, ECCM-4, qui se tiendra a STUTTGART (R.F.A.) du 25 au 28 septembre 1990.

PREFACE

Dr A.R. BUNSELL President of

the European Association for Composite Materials

Following on the success of both its predecessors, this third European Congress on Composite Materials, ECCM-3, is shaped as a top level event.

The first meeting took place in BORDEAUX in September 1985 and the second, which was twinned with the Sixth International Conference on Composite Materials, ICCM-VI in LONDON in July 1987. Both meetings demonstrated clearly the breadth of research on composites which goes on in Europe and served to bring together researchers from many countries, within Europe and from across the world.

This third conference brings us back to BORDEAUX and in doing so underlines the commitment that the South West of France, centred on this city, has made to the development of composite materials. A combination of political will and foresight, as well as industrial activity, has created an infrastructure centred on BORDEAUX which has made the activities of the European Association for Composite Materials, E.A.C.M., and the organisation of those European conferences possible.

The wide range of subjects and the high quality of the papers which has been asured by the diligent work of the scientific committee involving the cooperation of specialists throughout Europe ensures that the proceedings of ECCM-3 will be a work of reference. Perhaps more importantly ECCM-3 is bringing together once again researchers from across Europe and the rest of the world and contributes to the creation of a large international composite community.

A future date to be noted already is the Fourth European Congress on Composite Materials (ECCM-4) to be held in STUTTGART (West-Germany) from 25th to 28th September 1990.

INTRODUCTION AUX THEMES DU CONGRES

P. LAMica President du

Comite Scientifique de ECCM-3

La preparation du programme scientifique pour ECCM-3 s'est deroulee selon un processus tres classique, semblable a celui de ECCM-1. Les resumes re9us ont ete classes, puis envoyes aux lecteurs pour appreciation, chaque resume etant examine par au moins deux personnes.

Les progres de l'Association Europeenne des Materiaux Composites entre ces deux conferences sont apparus nellement lors de celie operation : pres de vingt-cinq responsables scientifiques de toute l'Europe ont participe a celie evaluation. Signe des temps aussi, I'emploi generalise du telefax a perm is des echanges d'informations flu ides et rapides, tout en evitant les voyages grands consommateurs de temps. Au total, une procedure souple, malgre I'augmentation du nombre de communications re9ues, et largement europeanisees par rapport a ECCM-1 ou un evaluateur sur deux etait fran9ais. L'efficacite est due en grande partie a I'action du secretariat, que je veux remercier ici de nous avoir grandement facilite la tiiche.

Nous avons rencontre quelques problemes de choix entre les sessions de presentation orale et la session de presentation par posters. Quelques auteurs ont exprime nellement leur preference pour I'une ou I'autre forme d'expression, ce qui a oriente certaines decisions. Pour les autres, nous avons privilegie la coherence de chaque session orale autour de son theme, tout en essayant d'y placer les communications, a caractere plus large, susceptibles de provoquer une discussion generale.

Les sujets plus specialises font souvent I'objet de discussions beaucoup plus passionnantesdevant un panneau, avec tout Ie temps necessaire et en petit groupe. Nous les avons mis preferentiellement dans une grande seance posters.

Les resumes que nous avons re9us a I'appui des propOSitions de communication nous ont paru d'un excellent niveau d'ensemble. lis annoncent tres probablement des communications de qualite. Cela est de bon augure pour que des debats animes puissent s'engager et pour que chaque participant a ECCM-3 ait Ie sentiment d'elargir et d'approfondir sa connaissance des composites.

Mon voeu, au nom de tous ceux qui ont participe a I'elaboration de ce programme, est que la communaute europeenne des composites se sente plus forte et plus unie grace a cette conference.

INTRODUCTION TO THE CONFERENCE THEMES

P. LAMica President of the

Scienfitic Committee of ECCM-3

The preparation of the scientific program for ECCM-3 was undertaken according to a very conventional process, similar to that of ECCM-1. The abstracts received were classified and then sent to readers for evaluation, each abstract being examined by at least two people.

The progress achieved by the European Association for Composite Materials between these two conferences appeared clearly from this operation ; close on twenty-five scientific leaders from all over Europe took part in this assessment. As a further sign of the times, the widespread use of the Fax enabled smooth and fast exchanges of information, while avoiding time-consuming travel. All in all, a flexible procedure, despite the increase in the number of papers received, and widely Europeanized with respect to ECCM-1, where one assessor out of two was French. Efficiency is largely due to the work of the secretariat, and I should like to thank them here for having considerably assisted us in our task.

We came up against a few problems in making a choice between the oral presentation sessions, and the poster sessions. Some authors clearly expressed their preference for one or the other form of expression, which guided us in certain decisions. For the others, we have pride of place to the coherence of each oral session around its central theme, while at the same time trying to insert wider-ranging papers likely to give rise to a general discussion.

The more specialized subjects often lead to much more lively discussions around a poster presentation, with all the time required, and with a small group. We therefore opted to include them in a major poster session.

The abstracts received to back the proposals for papers appeared to us to be of a very high level, in general. They are very probably the heralds of high-quality presentations. That augurs well for lively discussions and to enable each participant of ECCM-3 to feel that his knowledge of composites has been widened and deepened.

In the name of all those who have taken part in the preparation of this program, I very much hope that the European composite community will feel stronger, and more united, thanks to this conference.

SOMMAIRE TABLE OF CONTENTS

SESSION CHAIRMEN ............................................................................... XXI

PLENARY PAPERS ................................................................................. .

POLYMERES I POLYMERS:

"Mecanismes et cinctiques de reticulation de systemes thermo­durcissables en presence de renfon - relations structures proprietes thermiques" M.P. GRENIER-LOUSTALOT, P. GRENIER ............................................... 35

"Modified bismaleimides for carbon fibre composites" P. KONIG, H. STENZENBERGER, M. HERZOG, W. ROMER ....................... 43

"Enhanced bonding of fiber reinforcements to thermoset resins" G. SUGERMAN, S.M. GABA YSON, W.E. CHITWOOD, SJ. MONTE ............ 5 I

"influence of the thickening agents and some external parame­ters to the formulation on the viscosimetric kinetics of the pre­impregnated polyester" A. VALEA PEREZ, M.L. GONZALEZ, I. MONDRAGON ............................. 57

FIBRES:

"The modulus of alumina fibres containing mesopores dependen­ce of orientation distribution" M.H. STACEY ............................................................................................. 65

"Statistical mechanical breakdown of single fibres and micro­composites using video microphotographic techniques" H.D. WAGNER, L.W. STEENBAKKERS...................................................... 71

"The strength of tungsten-cored silicon-carbide fibres and the influence of a polymer matrix" M.G. BADER, D.A. CLARKE............ .......................................................... 79

MATRICES CERAMIQUES / CERAMIC MATRIX:

"Fibre reinforced alumina ceramic composites by sol-gel processing" M. CHEN, P.F. JAMES, F.R. JONES, J.F. BAILEy....................................... 87

"Silicon carbo-nitride ceramic matrix composites by polymer pyrolysis" R. LUNDBERG, P. GOURSAT...................................................................... 93

"Composites a matrice ccramique, nouveaux materiaux a tres hautes performances" '\ J.e. CAVALIER, A LACOMBE, 1.M. ROUGES ........................................... 99

"Interface characterization in ceramic matrix composites fabricated using FCYI techniques" R.A LOWDEN, D.P. STINTON, T.M. BESMAN ...................................... ;.... NC

"Resistance to crack growth in fibre reinforced cement : effect of fibre properties" L. DESCHRYVER, AR. BUNSELL, A. LE FLOC'H........................................ 111

"Thermo-mechanical characterization of ceramic composites made of a las glass-ceramic matrix reinforced with silicon carbide (nicalon) fibers" E. MENESSIER, A GUETTE, R. PAILLER, R. NASLAIN, L. RABARDEL, B. HOSTEN, T. MACKE, P. LESPADE ......................................................... 121

"Interface characterization by transmission electron micros­copy and Auger electron spectroscopy in tough SiC fiber (nicalon)-SiC matrix composite with a boron nitride interphase" O. DUGNE, S. PROUHET, A. GUETTE, R. NASLAIN, J. SEVELY............... 129

MATRICES METALLIQUES / METAL MATRIX:

"Cast fibre reinforced aluminium alloy microstructures" R. TRUMPER, V. SCOTT............................................................................. 139

"The influence of thermal cycling on the properties of Si3n4 whisker reinforced aluminium alloy composites" Y. NISHIDA, M.H. MASARU, M.Y. NAKANISHI .................................... 145

"Fracture of Al-SiCw metal matrix composites" D. CHAMBOLLE, D. BAPTISTE, P. BOMPARD .......................................... 151

"Physicochemical aspects of the K2ZrF6 process allowing the spontaneous infiltration of SiC (or C) preforms by liquid aluminium" S. SCHAMM, J.P. ROCHER, R. NASLAIN ................................................. 157

"Mechanisms for mechanical property enhancement of fibre reinforced metals using a hybrid technique" E.A. FEEST, RM.K. YOUNG, S.I. YAMADA, S.I. TOWATA ...................... 165

"Powder metall urgical prod uction of whisker reinforced magnesium" K.U. KAINER, J. SCHRODER, B.L. MORDIKE .......................................... 171

"Vibratory orientation of short reinforcing fibers in metal matrix composites"

B. SHPIGLER .................... .... ...... .... ....... .............. ...................................... 177

"Titanium matrix composites reinforced by C.V.D. filaments: a review of their thermo-mechanical capabilities" Y. LE PETIT-CORPS, T. MACKE, R. PAILLER, I.M. QUENISSET ............ 185

"A comparative study of thc mechanical behaviour of zinc reinforced by stainless steel filaments manufactured via two different processes" A. MADRONERO, M. PRENSA MARTINEZ-SANTOS ................................ 193

"Creep rupture of 1100 series A I/SiC particulate MMC'S" S. PICKARD, B. DERBy............................................................................. 199

"Microstructural stability of fibrous composites based on magnesium-lithium alloys" M. WARWICK, RT.W. CLYNE .................................................................. 205

"Microstructural development and mechanical behaviour of SiC whisker-reinforced Mg-Li alloys" 1. MASON, R.T.W. CLYNE ......................................................................... 213

"Particle reinforced magnesium alloys" 1. SCHRODER, K.U. KAINER, B.L. MORDIKE .......................................... 221

"Heat-treatment effects in Ii - alumina fibre reinforced alumi­nium alloy 6061" C. FRIEND, RYOUNG. I. HORSFALL ....................................................... 227

"Hot working behavior of discontinuous SiC/AI composites obtained by rheocasting" B. COUTAND, F. G1ROT, Y. LE PETIT-CORPS, J.M. QUENISSET .............. 233

INTERFACES:

"Etude de I'interface l"ibrc-ciastomcrc dans les composites a monofilament" M. NARDIN, E.M. ASLOUN, M. BROGL Y, J. SCHULTZ ........................... 243

"Greffage electrochimique de fonctions aminees en surface de fibres de carbone : cffet sur la tenacite d'un composite carbone­epoxy" B. BARBIER, M. VILLATTE, G. DESARMOT ............................................ 249

"Etude des proprietes de surfacc dc fibres aramides par chroma­tographic gazcusc inverse : crfct de divers traitements" J. SCHULTZ, L. LAVIELLE, A. BRUNERO ................................................ NC

"Elaboration en continu d'un depot mince de carbure refrae­tairc en surfacc des fibres de carbone : caracterisation de la fibre C/SiC" H. VINCENT, C. VINCENT, J.L. PONTHENIER, H. MOURICHOUX, J. BOUIX .......... ................ .......... ......... ........................................................ 257

"Sputter dcposition of diffusion barrier coatings on SiC mono­filaments for usc in Ti-based compositcs" R. KIESCHKE, R. SOMEKH, T.W. CL YNE ................................................. 265

"Sims analysis of SiC coated and uncoated nicalon fibers" M. LANCIN, J.S. BOUR .............................................................................. 273

"The effect of surface treatment on the interfacial strength of corrosion rcsistant glass fibres in a vinylester resin" F. JONES, D. PAWSON ................................................................................ 279

"Effeqs of matrix microstructurc changes after annealing on fracture properties of polypropylcnc/glass fibres injection molded composites" J. STEIDL, Z. KORINEK, V. ZILVAR......................................................... 287

"Compatibilite chimique entre Ie magnesium et les fibres de carbone" I.C. VIALA, P. FORTIER, G. CLA VEYROLAS, H. VINCENT, I. BOUIX..... 293

CONCEPTION ET CALCUL I DESIGN AND ANALYSIS:

"Etude de la fissuration d'un materiau composite verre-epoxyde unidirectionnel sollicite en traction" G. PLUVINAGE, C. SCHMITT, M. ABISROR ............................................. 301

"Design and analysis of orthotropic composite materials through a mixed mode cohesive crack simulation" S. VALENTE, A. CARPINTERI .................................................................. 309

"Contact behaviours of laminated composite thin shells and a rigid ball" LJ. LEE, C.K. PAl, L.c. SHiAU ................................................................. 315

"Design and analysis of statically and dynamically loaded composite sandwich panels" E. RASKER ................. .... .... ...... .................................................... ............. NC

"Large deflection initial failure of laminated rectangular plates" G. TURVEY, M. OSMAN............................................................................. 321

"Post-buckling of flat stiffened graphite/epoxy panels under cyclic compression" Y. FROSTIG, A. SEGAL, I. SHEINMAN, T. WELLER ................................ 333

"Finite element analysis of composite panel flutter" L.C. SHlAU, D.H. TSAY, L.l. LEE.............................................................. 341

"Optimising the geometry of energy absorbing composite tubes with particular reference to rail vehicle application" I.F. KELLy............ ..... .... ...... ..................... ................................... ............. 347

"Stresses in the joint of an end fitting to a composite torque tube" H. BROWN, R. HAINES, T. JOHNS, 1. MURPHy....................................... 353

"Strength and response of composite plates containing an open hole and subjected to compressive loading" F.K. CHANG. L. LESSARD .......... ............................................................... 359

"Design analysis of web-core composite sandwich panels" A.F. JOHNSON. G.D. SIMS ......................................................................... NC

"Integrally woven sandwich-structures" K. DRECHSLER, I. BRANDT. F.I. ARENDTS ............................................. 365

ELABORATION I PROCESSING :

"A model for pressure bag technology" I. CRIVELLI-VISCONTI. A. LANGELLA .................................................. NC

"The development and application of the multi live-feed moulding process for the production of injection mouldings containing laminated and other specific fibre orientation distributions" P. ALLAN. M.I. BEVIS .............................................................................. 375

"Fabrication of fiber reinforced ceramic composites" K. NAKANO. A. KAMIYA, M. IWATA, K. OSHIMA ................................ 381

"Processing parameters influence on the morphology and mechanical properties of sheet moulding compounds" I. IMAZ. A. RUBIO, C. FORURIA, J.F. LICEAGA ..................................... 389

"The manufacture of ultra-lightweight large diameter compo­site pistons" P. MOBBS ................................................................................................... 395

"Design of domes by use of the filament winding technique" M. MARCHETTI. D. CUTOLO, G. DI VITA ................................................. 401

COMPORTEMENTS MECANIQUES I MECHANICAL PROPERTIES:

"Statistical inference about stress concentrations in fibre­matrix composi tes" L.c. WOLSTENHOOME, RL. SMITH, M.G. BADER ................................... 411

"A standard for interlaminar fracture testing of composites" P. DAVIES, A. ROULIN .... ...................... ......... .... .... .......... ............... ....... 419

"Impulse and random tests for the modal parameters evaluation for a CFR panel" P. GAUDENZI ............. . ......... ...... ............................................................ 425

"Effet des conditions d'claboration sur Ie comportement meca­nique, statique et dynamique de materiaux composites hautes performances a matrice thermoplastique semicristalline" e. VERDEAU, A.R. BUNSELL ................................................................. 431

"Mechanical strength properties for anisotropic composites" e.L.D. HUANG ...... ........................ ..... ........ .... ............................... ........ ..... 441

"High performance composites made of solid thermoplastic powder impregnated fiber bundles" K. FRIEDRICH, H. WITIICH, T. GOGEVA, S. FAKIROV........................... 445

"Effect of fibre volume fraction on tensile fatigue behaviour of UD glass/epoxy composite" I. PARTRIDGE, P. VIRLOUVET, l. CHUBB, P. CURTIS ........................... 451

"Matrix selection for GRP fatigue loaded structures" A. GUEMES, l.A. GLEZ-VECINO, M.A. CASTRILLO ................................. 457

FISSURATION I CRACK:

"Influence of the fibre-matrix interface on the matrix crack development in carbon-epoxy cross-ply laminates" J. IVENS, M. WEVERS, I. VERPOEST, P. de MEESTER ............................ 465

"Analysis of thick laminates using effective moduli" e.T. SUN, W.e. LIAO .................................................................................. 473

"Microfractography of carbon fibre-reinforced bismaleimides" G. MAIER, M. WOHLEKE, P. VETESNIK, J. KUNZ ....................... ............ 48 I

"Porosity in advanced composite materials : its evaluation and effects on performances" A. GUEMES, M.A. MOLINA-COBOS, R. GONZALEZ-DIAZ ....................... 487

"Matrix cracking in cross plied thermosetting and thermoplas­tic composites during monotonic tensile loading" R. VAN DAELE, I. VERPOEST, P. de MEESTER ....................................... 493

ENDOMMAGEMENT ET FATIGUE / DAMAGE AND FATIGUE:

"Damage development in carbon fibre reinforced composite laminates under compressive static and fatigue loading" K. SCHULTE, J.1. MASSON ........................................................................ 501

"Damage tolerance of carbon fibre reinforced plastic sandwich panels" K. LEVIN ...................... ...... .......... ............................................................. 509

"Static and fatigue fracture of composites in complex state of stress" D. PERREUX, C. OYTANA, D. VARCHON .................................................. 515

"Damage development in CFRP and its detection" R. AOKI, J. HEYDUCK ........... ................................................................... 521

"Stiffness changes during fatigue of angle-ply glass/polyester of high quality under very large number of cycles" S. ANDERSEN, H. LILHOL T ......... .......... ................................................... 529

"3D-fabrics for composite sandwich structures" I. VERPOEST, M. WEYERS, P. de MEESTER ............................................ 535

FLUAGE / CREEP:

"Non linear viscoelasticity applied for the study of durability of polymer matrix composites" A. CARDON, H.F. BRINSON, C.c. HIEL .................................................... 545

"Comportement au f1uage de strati fies polyester/verre E destines Ii des applications navales" A. LAGRANGE, R. JACQUEMET ................................................................ 551

VIEILLISSEMENT / AGEING:

"Thermal fatigue of carbon fibre/bismaleimide matrix composites" T. JENNINGS, D. ELMES, D. HULL ........................................................... 563

"The influence of temperature and moisture on cross-ply cracking in CFRP in terms of matrix fracture strain and inter­face strength" P.W.M. PETERS, S.1. ANDERSEN .............................................................. 571

"Artificial ageing of fibre reinforced composite materials three stage method" N. MARKS, A. DALZIEL ........................................................................... 587

"Aspects of the thermal degradation of PMR-15 based composites" F. JONES, Z. XIANG ... ..................... .................................................... 595

"The hygromechanical degradation of aramid-epoxy composites" W. JANSSENS, I. VERPOEST, L. DOXSEE ............................................... '" 603

"Fatigue behaviour of GFRP some considerations about inter-f ac e s" L. VINCENT, L. FIORE, P FOURNIER ..................................................... 609

"Systematic fretting wear and fretting fatigue studies on carbon fibre/epoxy laminates" O. JACOBS, K. FRIEDRICH, K. SCHULTE .................................................. 615

"Influence of moisture on the compression behaviour of com posi tes" G. ZIEGMANN ....... .

"Moisture diffusion into two-phase matrix resins for fibre composites" F. JONES, P. JACOBS ................................................................................ .

"Influence du vieillissement sur Ie comportement au perl age de tubes verre-rcsine"

621

627

I. GHORBEL, D. VALENTIN. M.e YRIEX, J. GRATTIER ......................... 635

"Moisture absorption in n lIencc on thc mechanical properties of carbon/epoxy composites" I. MONDRAGON. J. IMAZ. A. RUBIO. A. VALEA ................................... 643

FRACTURE:

"The influence of fiber bundle reinforcement on the fracture mechanical behaviour of polycarbonate and epoxy" J. KREY, K. FRIEDRICH, K.H. SCHWALBE .............................................. 651

"Prediction of impact behaviour concerning CFRP laminates" A. RUBIO, 1.1. IMAZ, I.F. L1CEAGA ......................................................... NC

"Trigger mechanisms in energy absorbing glasscloth-epoxy tubes" I. SIGALAS, M. KUMOSA, D. HULL ......................................................... 657

HYBRIDES / HYBRIDS:

"The effect of agglomeration and the residual stress state on the performance of graded particulate hybrid glass fibre composites" F. JONES, S. AHMED ............. .................. ................................................... 665

"Compressive behaviour of unidirectional glass/carbon hybrid laminates" G. KRETSIS,F. MATIHEWS, 1. MORTON, G. DAVIES .............................. 671

"Fatigue of hybrid composites" B. HARRIS, T. ADAM, H. REITER ............................................................ 677

MODELISATION, SINGULARITES / MODELING:

"An analytical investigation on the thermally induced response of composites in the absence of thermal equilibrium" J. FLORIO Jr, J.B. HENDERSON, F.L. TEST ............................................... 687

"A tentative interpretation of the CFRP mechanical characte­ristics based upon the fibre/matrix relations at the interface. A case study" J.P. FAVRE, G. DESARMOT, 1. HOG NAT, 1. ROUCHON ............................. 693

"Endommagement en compression et en traction autour d'un trou d'un materiau composite carbone/epoxy" D. LAI, C. BATHIAS .................................................................................. 699

"The influence of specimen geometry and test conditions on the tensile and fracture mechanics properties of GRP" A.V. LIMA, J.A.O. SIMOES, A.1.M. FERREIRA, A.T. MARQUES, P.M.S.T. de CASTRO .................................................................................. 705

"Semi-empirical modelling of stress rupture data on glass reinforced plastics" 1. SILLWOOD, 1. A VESTON ........................................................................ 713

"A model of laminated composite plates assuring the continuity of displacements and transverse shear stresses" M. TOURATIER, Q. LlU, P. LORy............................................................. 721

"Free-edge stress singularity computation" P. DESTUYNDER, Y. OUSSET .......... ........ ................................................... 725

"Evaluation d'un nouvel clement fini pour I'analyse statique ou dynamique des plaques composites" P. LARDEUR, J.L. BATOZ .......................................................................... 733

CISAILLEMENT I SHEAR:

"An experimental-analytical investigation of intralaminar shear properties of unidirectional CFRP" W. BROUGHTON, M. KUMOSA, D.HDLL................................................... 741

"On the end notched flexure (ENF) test for the mode II interla­minar toughness of continous fibre reinforced composites" M. DAVIES, D.R. MOORE ... ....................................... .................. .............. NC

"Interlaminar fracture testing of carbon fibre/peek composites. Validity and applications" P. DAVIES, W. CANTWELL, H. RICHARD, C. MOULIN, H.H. KAUSCH... 747

METHODES NON DESTRUCTIVES I NON DESTRUCTIVE TECHNIQUES:

"La tomodensitometrie : methode non-destructive efficace d'observation des endommagements dans les materiaux composites" C. LAMBERT-CAMPELLO, C. BATHIAS ................................................... NC

"Measuring strain in carbon fibre composite laminates using the Raman optomechanical strain gauge" F. UNDERWOOD, D. SHARPE, D. BATCHELDER ....................................... 759

"Raman optomechanical studies on fibres and composites" G. GALIOTIS, H. JAHANKHANI, I. MELANITIS, D. BATCHELDER........ 765

"Developpement de methodes de controle par emission acous­tique des structures composites" C. HERVE, M. CHERFAOUI, M. TRUCHON, X. DUFOUR ........................... 771

"Cure characteristic determination using microelectronic dielectric sensors" D. DAY, H. LEE, K. RUSSELL, J. WHITESIDE............................................ 779

"The use of laser moire interferometry in the study of deforma­tion fields in composites and adhesives" R. DAVIDSON ............................................................................................ 785

"NDE of thick GFRP composites through ultrasonic waveform detection" R. TETI, G. CAPRINO ................................................................................ 793

CHOC / IMPACT:

"Inertial effects in twin skinned GRP laminates subjected to impact loading in a three point bend configuration" R.A.W. MINES, C.M. WORRALL, G. GIBSON ........................................... 803

"The effect of crystallinity on the impact properties of advanced thermoplastic composites" D. LEICY, P.J. HOGG .................................................................................. 809

Communication parvenue hors delai / late paper:

"Endommagement de structures tubulaires composites so us sollicitations dynamiques" P. HAMELIN, C. BURTIN.......................................................................... 819

SESSION CHAIRMEN

• POLYMERESIPOLYMERS Pr Dr K. FRIEDRICH, Technische Universitat Hamburg-Harburg

• FmRES Dr A.R. BUNSELL, Ecole Nationale Superieure des Mines de Paris

• MATRICES CERAMIQUES/CERAMIC MATRIX Mr B. BROQUERE, S.E.P. Pr B. HARRIS, University of Bath

• MATRICES MET ALLIQUESIMET AL MATRIX Pr M.G. BADER, University of Surrey Dr K. SCHULTE, D.F.V.L.R.

• INTERFACES Dr A.K. DHINGRA, E.!. Du Pont de Nemours Pr R. NASLAIN, Laboratoire des Composites Thermostructuraux

• CONCEPTION ET CALCUL/DESIGN AND ANALYSIS Pr F.L. MATTHEWS, Imperial College of Science and Technology Pr S.W. TSAI, US Air Force Dr lA.N. SCO'IT, Shell Laboratorium

• ELABORATIONIPROCESSING Dr M. KELLERER, M.B.B.

• COMPORTEMENTS MECANIQUES/MECHANICAL PROPERTIES Mr Th. JOHANNES SON , Institute of Technology of Linkoping

• FISSURATION/CRACKS Dr D.C. PHILLIPS, A.E.R.E. Harwell

• ENDOMMAGEMENT ET FATIGUE/DAMAGE AND FATIGUE Mr N. SPRECHER, Owens-Coming Fiberglas Europe

• FLUAGE/CREEP Dr A.R. BUNSELL, Ecole Nationale Superieure des Mines de Paris

• VIEILLISSEMENT/AGEING Dr H. Lll..HOLT, Riso National Laboratory DrG. GRUNINGER, D.F.V.L.R.

• FRACTURE Dr G. Di DRUSCO, Montedison

• HYBRIDESIHYBRIDS Pr I. CRIVELLI-VISCONTI, Universita di Napoli

• MODELISATION, SINGULARITESIMODELING Pr J.C. SEFERIS, University of Washington

• CISAILLEMENT/SHEAR Dr I. VERPOEST, Katholieke Universiteit Leuven

• METHODES NON DESTRUCTIVESINON DESTRUCTIVE TECHNIQUES Dr A. SA V AOORI, Enichem

·CHOCIIMPACT Mr P. LAMICQ, S.E.P.

PLENARY PAPERS

• "THE STRUCTURE AND PROPERTIES OF ARAMID FIBRES" 3 Dr M.G. NORTIIOLT, Akzo Research, Arnhem, The Netherlands

· "STATE OF THE ART AND FUTURE PROSPECTS FOR HIGH TEMPERATURE COMPOSITES" 5 Mr J.E JAMET, A6rospatiale Aquitaine, Saint M6dard en Jalles, France

· "APPLICATIONS OF TRANSMISSION ELECTRON MICROSCOPY FOR THE STUDY OF COMPOSITES (CARBONS-SiC). RELA TIONSIDP WITH MECHANICAL PROPERTIES" 15 Mme A. OBERLIN, Laboratoire Marcel Mathieu, Pau, France

• "ADVANCED MATERIALS TRENDS" 21 Dr J. BUCHANAN, BP Chemicals, London, Great-Britain

• "LIMITS TO TODAYS COMPOSITES· CHANCES FOR TOMORROWS DEVELOPMENTS" N.C. Dr H. KELLERER, M.B.B., Munich, West-Germany

ABSTRACT

THE STRUCTURE AND PROPERTIES OF ARAMID FIBRES

M. NORTHOLT

Akzo Research Laboratories P.O. Box 9300 - 6800 SB ARNHEM - The Netherlands

Aramid fibres are very tough and strong polymer materials. The modulus of elasticity of these fibres can range from 60 to 140 GPa, whereas the strength reaches a value of 4.5 GPa at 2.5 cm gauge length with an elongation at break of more than 4%. The origin for these properties is found in the intermolecular interactions resulting in a semi­rigid conformation of the poly(p-phenylene terephthalamide) chain. X-ray and electron diffraction studies have demonstrated that the structure is paracrystalline. On the basis of a single-phase structural model the elastic properties of this fibre have been fully explained.

Another outstanding property of the aramid fibres is the low rate for creep and stress relaxation. Presumably this is due to the semi-rigid nature of the chains, the high crystallinity of the structure and the hydrogen bonding between the chains. For a better understanding of the visco­elasticity of the fibre it is necessary to know the struGtural phenomena that happen during creep and stress relaxation. A useful tool is provided by the dynamic compliance measurement, because in well-oriented fibres it is linear related to the second moment of the chain orientation distribution. Experiments have shown that during creep and stress relaxation a progressive contraction of this distribution takes place, ·which is caused by shear deformation of the crystallites. These results have lead to a further development of the series model, which now incorporates viscoelasticity.

4

REFERENCES

1 - "On the crystal and molecular structure of poly(p-phenylene terephthalamide) fibres", M.G. Northolt and J.J. van Aartsen, J.Pol.Sci. Letters Ed. 11, 333 (1973).

2 - "Direct observation of structure in high-modulus fibres", M.G. Dobb, D.J. Johnson and B.P. Saville, J.Pol.Sci. Symp. Ed. 58, 237 (1977).

3 - "Aramids bridging the gap between ductile and brittle reinforcing fibres", M.G. Northolt in "Recent advances in liquid crystalline polymers", Elsevier Appl. Sci. Series, Chapter 20, 299-310 (1985).

4 - "Elastic extension of an oriented crystalline fibre", M.G. Northolt and R. van der Hout, Polymer 26, 310 (1985).

5 - "Viscoelasticity of aramid fibres", M.G. Northolt, J.H. Kampschreur and S. van der Zwaag, Proc. Rolduc Polymer Meeting 1988, to be published in "Integration of Fundamental Polymer Science and Technology" by Elsevier Applied Science in 1989.

ABSTRACT

ST ATE OF THE ART AND FUTURE PROSPECTS FOR

HIGH TEMPERATURE COMPOSITES

J.F. JAMET

A(Hospatia/e Aquitaine Centre Technique d'/ssac

33165 ST MEDARD EN JALLES- France

The interest and activity in the carbon and ceramic composites field is the result of their new potential ities and increasing promises. These composites meet the requirements of "thermostabil ity". reI iability and specific performances mainly imposed by the new projects of Aeronautical and Spacial Systems.

Without being the only reI iabil ity parameters. thoughness. damage stress levels and thermochemical resistance of these new materials are now determinative for futher developments and uses.

After an analysis of the microstructural rupture of these composites and of the evolution of models describing this rupture. this paper develops breifly the main aspects of carbon composites protection and ceramic composites manufacturin9 in order to emphasize: -the rules governing the relations between microstructure and rupture. -the consequences upon architecture. processing and using methods

1- INTRODUCTICIi

Carbon and ceramic composites can be considered as a new class of thermostructural materials on which large aerospacial projects are based. They concern the STP for space vehicules I ike HERMES and new generation of propellers. They are also taken into consideration for their dimensional stabil ity in space applications such as mirrors and structures. In addition. they are able to solve various problems relating to high temperature appl ications in which thermostructural (aeronautical propulsion) and phYsical properties

6

(electromagnetic) are needed. These uses mainly imply I ightness. high toughness and thermal

stabil ity, thermal cycl ing and damage resistance. sometimes, high emissivity and low catal icity. often large dimension parts and always reproducibil ity.

The development effort from which the

of such materials impl ied a large research state of the art will be described with the

present I imits of their using and future tendencies. On the thermomechanical point of view, the development of

long fibers ceramic composites is basically related to the previous development of carbon composites. These last composites were designed to fit military reentry and rockets nozzles environments (very high temperature and short duration).

Their manufacturing processes were basically reused for ceramic composites, that is to say: -multidirectional weaving. -CVI and liquid densification.

Their rupture behavior can be strictly compared with those of ceramic composites. Nevertheless, they represent a simpler case because carbon fibers have a higher thermal stab iIi ty compar"ed '" i th present ceramic fibers leading usually to drastic damages during processes.

The second analogy between carbon and ceramic composites concerns reusabilitY.These two materials are presently competitive because the first one can be effectively protected against oxidation and because the second one uses more and more knowledges developed for the first one.

So, a second generation of high temperature composites is now in progress. It enforces rigourous approaches specially to connect rupture behavior, microstructures, processes and reactivity with environment together. This implies a multipurpose approach which enter mechanic, thermochemistry, sol id physics, microanalysis, •.•

2- FRACTURE MECHANIC AND RELIABILITY OF HIGH TEMPERATURE COMPOSITES

By 1980, significant progress occured in terms of toughness with the advent of ceramic composites. This evolution is connected with a better knowledge of the toughening mechanisms, the development of new ceramic fibers and new densification methods.

The main difference between ceramic (or carbon> composites and other one is based on the relative positions of ultimate elongations of their matrices and of their reinforcements. Their 1 inear stress-strain diagrams illustrate this difference (figure I>. They introduce the critical stress ( ) of ceramic composites at which the ultimate strain of the matrix is reached. So, an increasing of the volume fraction (Vf) and the elastic modulus (Ef) of the reinforcement increases this critical stress. The ACK theory /1/ has explained the multicracking mechanism occuring at this critical stress.

More recently, a statistical model of the rupture /2/ gave a satisfying explanation of experimental results revealed by real stress-strain diagrams of ceramic composites (see figure 2) /3/: -increasing of the ultimate elongation of the matrix with Vf -progressive multicracking,

7

-hysteresis loops after the end of the multicracking. For example, such statistical stress-strain diagrams are

presented figures 3 and 4 /2/. They respecti~ely emphasize the influence of Vf on the first crack stress and of on the multicracking e~olution. This approach introduces flaw and dimension effects at the microstructural le~el. So, it is clear that the high toughness of ceramic (or carbon> composites requires simultaneously: -a ultimate strains of the reinforcements higher than those of matrices, -interfacial conditions which allow sliding when cracks reach fibers, -matrix multicracking and fiber bridging.

Consequently, beyond the critical strain, the ~ulnerability of such composites increases progressi~ely. Ne~ertheless, it is ad~isable to discriminate between composites for which the ratio EfVf/8mVm is below and abo~e unity. In the first case, the potential energy in the matrix before cracks initiation is always higher than those in the fibers. In such a case. the energy a~ailable in the matrix is generally enough to de~elop a total cracks propagation in the matrix. Such a beha~ior has been identified for instance in SiC/SiC and in some Carbon/SiC composites. When this ratio is abo~e unity. the energy a~ailable in matrix after initiation can be unable to achieve the cracks propagation. This is particularly true when the matrix shows a pronounced subcritical beha~ior. SiC/LAS and carbon/glass composites present this beha~ior when Vf is sufficiently large. In other words, the le~el of ~ulnerabil ity of ceramic composites introduced by the multicracking depends on the components properties.

3- PROTECTED CARBON-CARBON COMPOSITES

Presently, there are se~eral ways to protect carbon against oxidation. They depend upon the nature of the carbon (glassy carbon, polycristall ine graphite, pyrolytic carbon •••. ) and mainly upon the temperature range of use. Beyond 1000 C, the most common method is based on silicon compounds coatings chemically compatible with carbon. This compatibil ity of sil icon compounds with the carbon is obviously guaranteed with the covalent sil icon carbide. Ne~ertheless, it presents intrinsic I imitations in connection with the thermochemistry of its reactivity with Si02. On an other hand. to get a real efficiency of such a coating under severe thermochemical environments, many technical difficulties have to be simultaneously solved, such as: -chemical bonding between carbon and sil icon carbide coating, -thermomechanical adequation between the composite surface and the coating, -gas-tightness allover the exposed surfaces, -in depth protection to a~oid a rapid oxidati~e collapse.

The first method consists in sil iciurizing carbons of the composite surface. This treatment ~oncerns a depth of se~eral hundred microns which realizes a porous SiC-bed. Two main sil iciurization methods are used: they are based upon the following reactions with carbon:

Si + C ----} SiC up to the silicon melting temperature

8

SiO + 2C ----) SiC + CO

Nevertheless, to meet the previous conditions, the SiC-bed has to present several specifications: l-homogenious and low Young modulus, 2-progressivityof the transition between carbon and SiC 3-realization of a first protection able to avoid rapid oxidative collapse, 4-realization of the conditions to receive an over coating able to guarantee the gas-tightness.

Generally, such a bed can't maintain a full protection during a long time. It reduces considerably the oxidation rate of the composite and may be consider~j as an in depth protection. To obtain the fu~l protection during a long time inside the limits of thermodynamical stability of such a system, the previous bed has to receive an oyer-coating. Here is the more critical phase because the choice of this over-coating is connected with several parameters subordinated to: -mean elastic properties of the carbon-carbon (mean Young modulus), -loading mode (tensile, compression, shearing), -thermal and mechanical stresses, -temperature and thermal shock, etc •••

In other words, two main approaches are used to choose the technology of this overcoating: ~ell infiltrated coating compatible with SiC, -Glass coatings at lower temperature.

The following mass loss rates diagrams (fig.5) compares the behaviors of a 2,5D carbon-carbon composite from Aerospatiale: -without treatement, -with siliconization treatment, -with siliconization treatment + glass coating. They have been measured in air at the atmospheric pressure.

Aerospatiale used a similar technology to protect its prototype of leading edge during the 83 phase of HERMES project. The thermomechanical tests effected in air at the atmopsheric pressure in the solar furnace of Almeria. These tests consisted in eight thermal cycles up to 1550 C during 30 minutes with specified loadings. After these tests, very low mass losses were measured and the compliance of th! leading edge was unchanged.

Such a result on a real structure emphasizes the potentialities of this way for which new processing are in progress to maintain the performances in more drastic environments.

4- THERMOMECHANICAL POTENTIALITIES OF CERAMIC COMPOSITES

High temperature limitations of ceramic composites are mainly related to the intrinsic stabil ity and reactivity of fibers. Secondly, to maintain toughness, the interface has to be protected or not modified specially after multicraking. This basic conditions of composites survival in the high temeprature range are presently reyeiwed through the thermomechanical properties of auailable ceramic fibers and some ceramic composites.

9

Properties of ceramic fibers

Among current commercial ceramic and carbon fibers presented on the following strength-Young modulus diagram (fig.6) few of them are suitable to receive a ceramic matrix because processes impose temperature up to 1000 C. Presently, most of the ceramic composite research and development programs concern essentially: -High strength or high modulus carbon fibers (d = 1.7 to 1.9) -Silicon carbide fibers (Nicalon, Tyranno or CVD-AVCO) (d = 2.5 to 3.2) -accessorily, mull ite or alumina fibers ( 3H,Sumitomo, Du Pont) (d = 3.1 t04) Their strengths are located between 2000 and 3300 HPa, their modulus between 200 and 450 GPa.

Whiskers, in term of structural stability, strength and modulus are theoritically remarkable. For example, Los Alamos SiC whiskers overtake a strength of 21 GPa and a modulus of 700 GPa. When their manufacturing will be commercially developed to get sufficient length, adequate packaging, low price and availibility, they will compete favorably with the best previous long fibers in ceramic composites processes if security problems are solved. Presently, very encouraging results have been obtained for composites with various matrices 14,5,61 and SiC whiskers. For example SiAION reinforced SiC whiskers manufactured by Aerospatiale, reaches a flexural strength of 450HPa with a fracture toughness of 11 HPaVm, and the creep rate at 1200 C is redu~ed by several orders of magnetude. Such processes using hot-pressing or hipping are suitable to make small parts for joining components (Hermes Program) and in the future they are considered for engine appl ications (turbine blades). Nevertheless, this class of process is not really fited to manufacture large structures.

Now, ultimate strengths evolutions of previous continuous fibers with temperature are very different (fig.7): if carbon fibers are quasi insensitive to the temperature in reducing atmosphere, ceramic fibers begin to lose their strength from 900 C but maintain their modulus up to this temperature. SiC-CVD fibers are presently the best in spite of their microstructural change, but they have two main disavantages: their large diameter (140um) and their subsquent difficult handiness. Now, SiC Nicelon fibers present important advantages (low diameter, weavingness) and are presently strongly investigated in the ceramic composite area. Nevertheless, their sensitivity to oxidation is well known /7,81 and begin beyond 600 C. In terms of high temperature stability, recent investigations have shown that they are largely perfectible /9/. So, in reducing atmosphere, their strength can be maintained'(90-93X) up to 1350 C, after long time heat treatments (several hours). Nevertheless, during such treatments, they produce few SiD, a very thin carbon coating (100A) and a I ight -SiC grain growth. Now, new similar fibers, produced from organometall ic precursor, are presently in development (Ube, Avco. Rhone Poulenc, Dow Corning HPDZ,HPZ and HPS). Their improvment is essentially based on titanium (Ube) or nitrogen (R.P. and D.C.) additions to reduce grain gr~th at high temperature. Simultaneously, a large effort is performed to reduce oxygen content (under sil ica form) which is 1 iable to high temperature reactions

10

either with free carbon or amorphous SiC 110/. It is too early to know their high temperature performances and their impact towards effective ceramic matrix reinforcement.

Oxides fibers are very sensitive to grain growth and creep when a glassy phase is present. They have been largely investigated in recent research programs. Precautionary measures have to be taken specially in oxide matrix composites, because of fiber-matrix interactions of particular importance. Like SiC fibers manufacturers, 3M and Du Pont develop presently new oxide fibers (Nextel 480 and PRD-166). The first one is mullite (+ ZI. 8203), the second is -Alumina with Y203 partially stabil ized Zr02 as a second phase 111/. All these modifications were made to improve their strength retention after exposure at high temperature through a presumable grain growth inhibition. Here too, it is too early to know their high temperature behaviors in ceramic composites. Nevertheless, recognizing the need of more refractory ceramic fibers, all these suppliers are working toward improved products.

High temperature properties of ceramic composites beyond 1000 C

For continuous fibers reinforcement, several manufacturing processes have been used with carbide, sil icate, oxycarbonitride and nitride matrices. Descriptions of specific combinations of fibers, matrices and processing conditions have been recently reviewed 111,12,13/. High temperature mechanical properties data of the corresponding composites are I imited. However, with SiC/glass ceramic composites several authors mention a rapid degradation occuring above 1000 C in air 1141 and transition in tensile failure mechanism to a brittle mode of failure occuring above 800 C in air /15/. So, among all these combinations, SiC/SiC /16/ seems to be the only long fibers ceramic composite able to sustain using beyond 1000 C.

The Nicalon 2D reinforced SiC-CVI composite (Cerasep 320) has a density of 2.75 and a critical stress of 100-IIOMPa (beginning of multicracking). Its initial tensile strength i's 210-23OMPa. The intrinsic oxidation resistance of such a composite is given by residual strengths after oxidation treatments under air (Table I). After these oxidations, the toughning mechanism is maintained but fibers degradation would be responsible for strengths decreasing. Their oxidation vulnerability after preloading is drastic above 105MPa. So to improve this composite, SEP has had to develop an oxidation protection which doesn't modify the toughning mechanism up to 1400 C during 10 hours (see figure 8).

In such conditions, Nicalon fibers are able to maintain ultimate strain up to O.T/'. This verifies the previous experimental potential ities of Nicalon fibers up to 1400 C if they are used in an adequate environment.

The carbon 2D reinforced is a thermostable system under strengths are quasi independent <Table 2>' It presents a large composi tes.

SiC-CVI composite (Sepcarbinox Ill) inert atmosphere. High temperature of the temperature up to 2000 C similarity with protected carbon

Under inert atmosphere, toughning is maintained up to 2000 C, but I ike the previous one, it needs an oxidation protection.

11

5- C()IICLUS I ()II

Presently, carbon and sil icon compounds (specially sil icon carbide and its sil ica protective layer) seem to be the two main materials able to maintain simultaneoustly high temperature stability and thermostructural properties of composites beyond 1000 C.

With· additional protection, Sic/SiC-CVI can be used safely over their critical stresses up to 1300 C.

Carbon composites with an efficient protection can be used beyond 1600 C. Silicon compounds protections systems have made great strides and this class of ceramic composites is really promising.

The inadequacy of SiC beyond 1600 C in low pressure and dissociated environments is now well establ ished and new protections methods of carbon are in progress for long duration purpose /17/.

SiC Nicalon for long fibers and SiC whiskers are the main ceramic reinforcements successfully used to day with various ceramic matrices. The first one may be used up to 1350 C with an efficient protection. The second one is very attractive for higher temperature and specially 1 ittle hot parts. Now, fibers suppl iers have recognized the need of more refractory products and are working towards to improved them. Nevertheless, carbon fibers are still the best reinforcement ,."hich is available to develop presently high temperature composites structures.

Now, up to 1000 C, carbon and SiC reinforced glass-ceramics present the best advantages in terms of manufacturing, performance and price. This is a promising way for a lot of appl ications in this temperature field.

REFERENCES

I. J.Aveston, G.Cooper, A. Kelly:Conf.Proceedings NPL (1971) 2. P.Peres, L.Anquez, J.F Jamet:Revue Phys.Appl.23 213-228(1988) 3. J.F.Jamet,D.Lewis,F.Y.Luh:ACS,Cocoa Beach 15-18 Jan.1984 4. P.F.Becher,G.C.Wei:J.Am.Ceram.Soc.67(12)C-267,C-26 (1984) 5. J.R. Porter and al :Am.Ceram.Soc.8ull. 66 (2) 343-47 (1987) 6. P.D.Shalek,J.Petrovic and al:Am.Ceram.Soc.Bull.65(2)351(1986) 7. T.Mah,N.L.Necht,D.E.McCullum:J.Mater.Sci .19(4)1191(1984) B. T.Clark,R.Arons,J.B.Stamatoff.Ceram.Eng.Sci .Proc.6(7-8)(1985) 9. J.F.Jamet, 12th Techcer,Rimini ,Oct.1987 10.T.Mah and al :Am.Ceram.Soc.Bull. 66 (22) 304-08 (1987) II.J.Cornie,Y.M.Chiang,D.R.Uhlmann.Am.Ceram.Soc.Bull.65(2) (1986) 12.L.J. Schider, J.J. Stigl ich.Am.Ceram.Soc.Bull .65(2)289-92(1986) 13.J.F.Jamet,L.Anquez,M.Parlier and al :A.A.A.F. 213 1231124 (1987) 14.J.J. Brennan, K.M. Prewo.J. Mater. Sci. 17 (8) 2371-83 (1982) 15.A.G.Evans and al :Proce.Vth Int.Conf.on Compo Mat.(1985) 16.J.F.Jamet,Institute of Physics Conference.Warwick, Sept.1987 17.A.P.Katz,R.J.Kerans:Am.Cer.Soc.Bul .67(8)1360(Aug.1988)

12

afu v, L -----acuL--­

I

._vm~ 'mu flu

fmu <'fu

8ri nl. mltrlx

CJcu • Em 'mu Vrn ~ E, 'mu V,

acu • em 'mu lit - v" ~ v, E,fem)

Fig. 1. Linear stress-strain diagram of unidirectional composites under tensile stress

1200 1080

360

840

720 600

4S0

~:~ 120 I o

o 2.4 6.4

'r

Fig. 3. Statistical stress-strain diagrams of unidirectional SiC/LAS composite: fiber content sensitivity

t fm rm

2.3 3.5 f x 10'

a 2 3 4

Fig. 2. Experimental stress­strain diagrams of an unidirec­tional SiC/LAS composite [1]

• IMP, I

Fiber rupture

11

50MPa 21

t= laMP, II , ,

t" 2MP, 41 .. 0 0.5 1 l 4 5 • 110",

Fill'. 4. Statistical stress-strain diagrams of SiC/SiC composite: sensitivity of the sliding strength at the interface

6

2

1.5

0.5

Specific H~ss Los! R~ t ps (KQlm2lH)

/

500

I

Non protected C~rbon Composite

Si I iconiud C~r bon Comp os i te

I ~ Siliconized / ~\ + Glass Coating

'--1000 1500 ( C)

Fig.5. Evolutions of specific m~ss loss rat.s of A.rospatial. Carbon Composit.s: -without prot.ction -si I iconized -sil iconiz.d + glass coating

In air at the atmosphere

2001SU- tMP.,

.ao

0.2

Pr,lO«1iN) 130M,",

+ IOhat 1'CIQ4C

Inmll compoSite

0.4

IOI'l'114OO"C

SiC - SiC CeI1lWP 320 I","o'lICwcll

0.' o. fig. 8. Thermomechanical beha­viour of protected Cerasep 320 after pre loading (130 MPa) and oxidation for Iv hours at 14000 C

13

EIG")

1 00 :zoo JOO &00 500

Fig. 6 . Mechanical properties of current commercial ceramic fibers

4 (GPu

200 400 600 800 1 000 1200

Fig. 7 . Thermomechanical properties of some ceramic fibers

Pre- Oxidation Oxidation Residual loading tempera- time strength

(IIPa) ture (OC) (h) (MPa)

70 1200 0.5 230 105 1200 0.5 221 130 1200 0.5 111 105 1100 100 148

Table 1. Residual strengths of Cerasep 320 after preloading and oxidation (*)

Sol icitation Temperature Strengths mode (OC) (MPa)

Tensile 20 350-400 .. 1200 450-500 " 1600 310-320

Flexural 20 450-500 .. 1300 650-700 " 2000 400-450 " 2200 250

Table 2. High temperature tensile and flexural strengths of Sepcarbinox 111 in an inert atmosphere (*)

ABSTRACT

APPLICATIONS OF TRANSMISSION ELECTRON MICROSCOPY FOR THE STUDY OF

COMPOSITES (CARBONS-SIC) RELATIONSHIP WITH MECHANICAL PROPERTIES

A. OBERLIN

Laboratoire Marcel Mathieu 2 avenue du President P. Angot - 64000 PAU - France

Peculiar modes of TEM allowing to restitute the three dimensional arrangement of elementary scattering domains (microstructure) of a given material are described. Applications to carbon and ceramic (SiC) fibres and matrices are given as examples. Mechanical and electrical properties related to microstructure are considered mainly for carbon fibres.

Les materiaux queUe qu' en soit la nature (fibres ou renfort, carbones ou ceramiques), presentent des proprietes physiques, parmi lesquelles les proprietes mecaniques et electriques, extremement dependantes de leur microstructure. Si ces materiaux ne sont pas macrocristallises, I 'arrangement dans I 'espace de leurs unites structurales de base est au mieux a l'echelle nanometrique ou micrometrique de sorte que la seule technique efficace pour les caracteriser est la microscopie electronique par transmission a haute resolution (MET).

La difficulte a surmonter pour etablir avec certitude la microstructure (c'est a dire en fait la microtexturel d'un materiau est de passer de l'image microscopique, qui est une projection orthogonale a deux dimensions d'un objet d'epaisseur non nuUe, a l'arrangement tridimensionnel. Pour ce faire, l'exploitation fine et complete de toutes les donnees cristallographiques fournies demoyennees par Ie MET est indispensable.

- un obj et periodique ne diffracte sous I' impact d' un faisceau d'electrons que lorsque ses plans reticulaires denses sont quasi paralleles au faisceau incident (loi de Bragg pour une longueur d'onde

16

petitesoit2dhle~).., avec).., (3.10-3 nm)6) coherent ou

tolerance sur desorientation

- plus l'u~ite structurale de" base (domaine ~ristallite) est petite ou mal organisee, plus la l'orientation de l'objet est grande. (10 0 et plus de tolerable pour un objet periodique)~

On peut alors imager dans un carbone toutes les distorsions des couches aromatiques individuelles dans les trois dimensions de l'espace.

On peut identifier et imager dans une fibre ceramique toutes les phases en presence et leur situation relative.

On peut identifier et imager dans un composite les interfaces ou interphases.

On peut, de la meme maniere, etudier la matrice. On peut enfin etudier Ie passage d'un precurseur au produit

fini : fibre ou matrice. Pour comprendre les principes de l'imagerie particuliere a employer

avec un TEM, ,on rappellera d'abord que la lentille objectif de l'appareil ramene les franges a l'infini que donne l'objet diffractant (diagramme de diffraction electronique) a son plan focal image. Ainsi, la formation de l'image est en relation de reciprocite parfaite avec Ie diagramme. A tout faisceau diffracte correspond une image ou la region diffractante est lumineuse sur un fond noir, tandis qu'au faisceau transmis sans diffraction correspond une image sur fond clair dont toutes les images lumineuses ont ete soustraites. La superposition de toutes ces images complementaires donne l'image proprement dite. Celle-ci prendra tout son interet lorsqu'on l'aura fi I tree en laissant se superposer une partie seulement des images portees par chaque faisceau. Par exemple, une imagerie en franges de reseau (F.R) correspond a une image d'interference. La trace des plans atomiques denses apparaltra sur I' ecran d' observation so us forme de franges. On pourra aussi filtrer l'image en ne recueillant que l'image lumineuse sur fond noir produite par un seul faisceau diffracte : imagerie en F.N.

- la condi don Q) de diffraction qui impose aux plans denses d' etre quasi paralleles au faisceau incident fait que l'image de ces plans, memes s'ils sont continus, disparait des qu'ils s'eloignent de leur position de Bragg.

- la condition (j), qui admet une erreur assez grande dans I 'orientation relative de l'objet periodique et du faisceau, aide a conserver l'image en F.R. et en F.N. de fa~on assez persistante pour rendre ces techniques utilisables dans un domaine d'autant plus large que l'objet est plus imparfait.

- c'est ainsi que pour les fibres de carbone (1) dont l'arrangement des couches aromatiques suit un modele en serviette plissee (fig. 1a et b) on a pu determiner, jusque dans Ie detail, la taille moyenne L , l'arrangement et les distorsions (rayons de courbure r l et 1Lt

parallele et perpendiculaire a l'axe de fibre). Dans les fibres haute resistance la couche parfaite est reduite a un diametre de 1 nm et l'element diffractant (unite structurale de base) ne comporte pas plus de 2 a 3 couches empi lees en desordre turbostratique. Les nappes distordues sont tres riches en defauts trappes aux frontieres des USB (fig. 1a) ~ est trop petit pour etre mesure. Dans les fibres haut

17

module obtenues a haute temperatures, les USB ont coalesce en couches de plus grand diametre (L > 20 nm) dont les rayons de courbures sont plus grands (3 < r t < 10 ~m) (B < l" < 250 nm).

La resistance a la rupture cro}! avec la cohesion laterale de la fibre c'est a dire avec sa compacite D t ou S =1 (11f t + l/rl). Lorsque deux nappes adj acentes se touchent, une soudu'fe peut s' etre formee par depart d'azote entre deux USB en contact. 6C est donc liee directement avec la densite de soudures (fig. 2a). De la me me maniere on a montre que E Ie module d'Young et la resistivite varient avec S et I (fig 2b et c).

_a A titre d'exemple on montrera dans une fibre "SiC" commerciale (2) 1 'association de microcristaux de SiC, de carbone libre aromatique et de silice amorphe dont I 'arrangement se modifie lors d'un depot CVD de SiC ou de carbone, pour aboutir a des interphases complexes (Si02, C ou SiC + C).

11 en resulte des modifications dans I' adhesion a la matrice SiC cristallisee, ce qui modi fie les proprietes mecaniques.

18

Fig. 1a - Modele de fibre haute tenacite

Fig. 'In - Modele de fibre 'hdut lliodule

19

O"c: GPa

3

2

1/S . . o 0.1 0.2 AlA

Fig. 2a - Courbe 0" c = f(1/S)

G Pa E

400

300

l --1

I I I La 0 1 • 1 2 3 4 5 10-3 A-

Fig. 2b - Cdurbe E = f(l/La)

20

p!2.m p

10

5

o 2 3 4

Fig. 2c - Courbe p= f(l/La)

REFERENCES

1 - Guigon M., Relation entre la microtexture et les

proprietes mecaniques et electriques des fibres de carbone

ex. Polyacrylonitrile, These d'Etat, 13.11.85, Compiegne.

2 - Maniette Y., Contribution a l'etude de phenomenes

d'interphases dans des composites de carbure de silicium,

These, 8.12.88, Pau.

ABSTRACT

ADV ANCED MATERIALS TRENDS

J. BUCHANAN

BP Chemicals Belgrave House, 76 Buckingham Palace Road

LONDON SW1 W OSU - England

Advanced materials is a new and growing business built upon substitution. Although materials scientists have the molecular understanding to design new materials, entry to the market is not guaranteed. Industry participants must face technological, economic and commercial challenges which can take between 20 and 40 years to overcome. The industry is also competitive and profitability is not assured. Nevertheless, significant progress has been made by polymer composites in aircraft 'structures, aero propulsion and marine applications and by ceramic materials in electronic and structural applications.

I - INTRODUCTION

Materials Technology is entering a new era of development as existing commodity materials are less able to meet the stringent demands of new applications. Materials scientists are developing the molecular understanding to enable us to design their successors to meet specific requirements. Through substitution, a family of new advanced materials have entered the market and are beginning to climb the product life cycle curve (Fig. 1).

With annual growth rates projected to exceed 10-15% over the next decade the development of these advanced materials presents a potentially attractive long term business opportunity which has been targeted by several major international corporations, such as BP.

22

II - INDUSTRY CHARACTERISTICS AND REQUIREMENTS

What are "Advanced Materials"? For the purposes of this paper they are formulations of metallic, organic and inorganic materials developed for specific structural or functional applications (Fig. 2). Prominent members of the family are engineering plastics, re-inforced plastics, advanced polymer composites, carbon/carbon composites, structural ceramics, electronic ceramics, ceramic composites, metal matrix composites and superalloys. These materials typically penetrate markets through offering higher performance to meet mechanical, environmental or electrical criteria.

Development and commercialisation of anyone of these materials takes money and patience. Indeed, it is conceivable that some of them will become themselves substituted before they reach the stage of commercialisation. For the winners, market development follows a similar and familiar pattern as it proceeds up the value chain (Fig. 3). For most materials this cycle can take 20-40 years (e.g. polyethylene, still maturing 1933-89) and involves three necessary inputs. The first is the fundamental scientific understanding of materials properties to enable judgement of the potential for any given line of development. The second is the technical appreciation of a chain of steps encompassing materials selection, processing, design and fabrication. And the third is the specific applications engineering necessary to tailor materials to meet customer requirements in a cost effective manner. All this must be underpinned by a conscious ongoing effort to educate designers and engineers to make optimum use of new materials in both new and current applications.

In following this path to commercialisation advanced materials first enter the market in applications where the ability to pay for enhanced performance is high, e.g. aerospace and defence. But as processing knowledge and design experience grows, the materials move into higher value volume applications and wider markets. Stimulating this process requires an understanding of the substitution economics which drive the demand for materials in the more traditional markets such as automotive, construction, industrial and electrical. Recognising these trends and requirements, major corporations with long term perspectives and strong commitments to technology development have entered the industry on a global scale. Their presence has elevated the level of competition and with that raises the spectre of an era of profitless growth.

III - OPPORTUNITIES

The world market today for advanced materials has been estimated at more than $25 billion. This is made up 50% each by structural and electrical usage and split evenly between polymeric and ceramic materials. By the turn of the century we forecast this market to grow to over $60 billion, almost two thirds of which will be structural materials, and around one third will be ceramics (Fig. 4).

This growth will come not only from increased usage of state of the art materials through a more comprehensive dialogue between the applications engineers of materials suppliers and major customers, but also from technical solutions to a number of problems which are restricting the use of these materials today. The general thrust of these te~hnical solutions will be in the enhancement of one or more of the following: strength,

23

stiffness, lightness, temperature resistance, corrosion resistance, wear resistance, impact resistance, thermal stability, electrical properties and optical properties, together with improvements in the processing of the materials to reduce labour costs, increase reliability and reproducibility and eliminate waste.

IV - SOME CONSTRAINTS TO DEVELOPMENT

Much R&D effort is currently being expended by government institutions, airframe manufactures, the automotive industry and materials suppliers towards these objectives but regardless of the outcome of these programmes, success in the market place, the acid test for all materials, will only come about when customers cost expectations as well as technical requirements are met. In this regard the introduction of new materials to the market is subject to a number of constraints. Firstly there is a strong likelihood that the new materials or process will be aiming to replace another which has a solid track record. Unless there is a crucial performance or cost advantage involved there will likely be a great reluctance to change. Furthermore there is also the possibility that the existing technology will be upgraded in reaction to the threat, witness the aluminium industry's move to superalloys and metal matrix composites in response to the perceived threat from novel materials. Then there is the problem imposed by user's capital equipment and infra­structure costs. There is always a tendency for companies with limited resources to try to ensure that any new material or technology is compatible with existing equipment and planned capital expenditure profiles. Finally there are the industry qualifications and legislative constraints associated with meeting performance and environmental standards. In some cases the lack of established user and environmental standards positively inhibits the introduction of the new technology as it

24

causes uncertainty and erodes user confidence. It is therefore implicit on us, the suppliers, to ensure adequate data on our products and the development of the legislative framework needed to support our industry.

v - TRENDS AND SUCCESSES

Let us now turn from the general industry characteristics and the concerns and problems facing suppliers and users to look at some of the trends in the industry today.

5.1 Polymer Composites in Aircraft Structures

Composite materials were initially used in aircraft to save weight and thereby improve overall operating efficiences. Their introduction began in non critical areas such as interiors and cosmetic features, using mainly glass re-inforcements. As customer confidence grew, high wear and structural applications were targeted resulting in the introduction of carbon and aramid re-inforcements and the application of composites into first secondary structures such as fairings, external pods, tanks and control surfaces which sub-contractors like BP Advanced Materials now produce on behalf of the aeroplane manufacturers, and finally into primary structures and skins. This extended use of composites in aircraft skins also makes feasible the introduction of EM shielding within the materials to impart "Stealth" characteristics to the airframe. Contrast the aircraft of the early 1970's to those of today as composites have penetrated from 0% by weight usage to as high as 45% in the latest generation of military fighters and in Beech Starship the first commercial aircraft with an all composite airframe and skin. Today the performance capability of composites is restricted not by stress limitations but by operating temperatures and impact resistance as aeroplanes fly faster and higher. The market pull is for stronger and stiffer fibres and tougher, higher temperature, solvent resistant matrices perhaps shifting the balance from thermoset to thermoplastic resin systems in the long term. Approved military programmes in both USA and Europe will ensure a growing market for some of these materials over the next 20 years.

5.2 Aero Propulsion

The aero propulsion story is similar to that for aircraft structures except for the need for a wider range of materials to cover a spectrum of operating temperatures. Polymer composites made their first appearance outside

the engine in the enclosing nacelle structures to reduce weight and provide better acoustics. They have subsequently moved into higher temperature and more structural applications; such as blocker doors which redirect the gas flow providing reverse thrust to assist braking the aircraft on landing. With resins such as PMR/15 now able to operate at temperatures of 300°C this trend is expected to continue. In the hotter parts of engines metals still reign supreme but materials such as carbon/carbon and ceramics offer potential for the future. Carbon/carbon is used extensively by BP's Hitco subsidiary in ablative applications where materials progressively erode in a controlled manner under high temperatures e.g. rocket motor exist nozzles, f1exsea1s which permit changes in direction of rocket motor thrust and heatshields for space launch vehicles and inter­continental missiles. Its major drawback is lack of stability in oxidising conditions and here lies another technology challenge for today's application engineers. Ceramic coatings are one answer which combines the structural strength of the substrate material with the temperature and corrosion resistance of the ceramic and brings about a marriage of two materials technologies -perhaps a trend for the future.

5.3 Polymer Composites in Marine Applications

The corrosion resistance, impact strength, lightweight and radar transparency of polymer composites have led to their use in marine applications. Whilst sill gaining acceptance with designers and marine architects, they have nevertheless been concerned in some striking developments. Early uses in the Royal Navy were for radar linear arrays and the gunshie1ds on frigates but in the USA, these materials are now used by Hitco for submarine bow domes whic"h at 12 m in diameter are the largest single autoclave moulded plastic structures made

25

in the world today. Extension of the use of new materials into a conservative industry such as shipbuilding is dependant upon applications engineering rather than development of new technology. It is a good example of the resistance to change which is inherent in many traditional markets.

Moving away from polymer composites and looking at markets other than aerospace and defence the development of new ceramic materials for electronic, industrial and automotive applications is another exciting area of the advanced materials business.

26

5.4 Electronic Ceramics

In electronics applications ceramics are competing with metals, semi conductors and polymers where they offer the prospect of enhanced corrosion resistance, dimensional stability, dielectric constant, thermal conductivity, hermetic seal, magnetic properties, electrical resistance and electro-optical properties. Already an $8 billion business based largely on alumina, barium titanate and iron oxides, electronic ceramics is looking towards increased performance and reliability in smaller units and wider applications for piezoelectric materials. New materials such as boron nitride, aluminium nitride, silicon carbide, various glasses and alloys are under development by BP's Carborundum subsidiary and others for packages, substrates, capacitors and piezoelectric devices. One major issue in the trend towards greater miniaturisation is the low product yield and this provides a technical challenge to process engineers in the opening of wider market opportunities.

5.5 Structural Ceramics

In structural applications ceramics offer extended resistance to wear, corrosion, elevated temperatures and thermal shock. Current applications are in automotive catalyst supports and pump seals where Carborundum is a market leader, spark plugs, cutting tools, nozzles and heat exchangers. Their main drawbacks are brittleness, difficulty to form into complex shapes and the low materials yields from current processing technology. The drive for increased toughness has started the development of ceramic matrix composites using particulate, whisker and continuous fibre re-inforcements. These show early promise in cutting tools.

Nevertheless, the growth prospects for ceramics in structural applications remains constrained today more by technology barriers in materials formulation and processing than by applications engineering - possibly the reverse of the situation in Polymers.

5.6 Metals

Finally, where the technical challenges remain beyond the capability of polymers or ceramics, there is still a major place in tomorrows world for metals. New processing technology such as superforming, new alloys such as aluminium-lithium and metal matrix composites offer the scope to stretch the current regime of metals and displace

the more costly new polymeric materials in applications where performance enhancement is not substantial enough to be cost effective.

VI. CONCLUSION

In summary, our business of advanced materials has assured future growth although it presents us with many imponderables and challenges, not least being making satisfactory returns. There are many contenders and whilst it is not possible to guess with any degree of certainty which materials will ultimately be winners, it is not clear that many wholly new materials will be discovered. Rather the emphasis over the next decade will be technology development and applications engineering to stretch existing materials and, through novel process engineering, improvement in the economics of their fabrication and use to widen available markets.

27

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Chairman: Pr Dr K. FRIEDRICH Technische Universitat Hamburg-Harburg

MECANISMES ET CINETIQUES DE RETICULATION DE SYSTEMES THERMODURCISSABLES

EN PRESENCE DE RENFORT RELATIONS STRUCTURES PROPRIETES THERMIQUES

M. F. GRENIER-LOUSTALOT. P. GRENIER

Centre National de /a Recherche Scientifique LCOP - UA 474 Avenue de I'Universite - 64000 PAU - France

ABSTRACT

Using a series of di- and tetra functional prepolymers we have shown the role of non epoxide chain extremities and impurities present in the prepolymers on the reaction mechanisms and kinetics and the final thermal properties of the materials.

From a mechanistic standpoint in the IOO-180·C temperature range, epoxide / primary amine reactions dominate at the beginning ; at the end of the reaction, epoxide / hydroxyl and epoxide / secondary amine reactions dominate. These reactions may be intermolecular (to form cross-linked net­works) or intramolecular to form rings, especially at the end' of the heating cycle.

On prepregs we were able to show the non-negligible role of the fibre in reaction kinetics but change reactions mechanisms only slightly.

INTRODUCTION

La caracterisation de matrices organiques pour materiaux composites necessite une connaissance precise de la structure des materiaux de base mais aussi de leur evolution lors de leur mise en oeuvre et la fabrication des composites. Plusieurs travaux de la Iitterature (1-4) ont montre I'impor­tance des impuretes et la presence de bout de chaine-fonctionnel non epoxyde dans les composes de base qui peuvent modifier la structure du reseau final.

A partir des resultats obtenus au laboratoire sur des composes modeles proches des prepolymeres industriels (5), nous avons etudie les mecanismes et les cinetiques de reaction a I'etat fondu de resine type TGDDM (tetra­glycidyl diamino diphenyl methane) commerciale et purifiee au laboratoire polycondensee avec du diamino diphenyl sulfone (DDS) afin de connaitre Ie role des impuretes et des bouts de chaines reactifs sur les cinetiques et la structure du reseau final.

36

A cette fin nous avons sur Ie prepolymere tetraepoxyde de depart iden­tifie les impuretes par synthese du produit et etudie d'une maniere precise un prepolymere commercial (LOPOX 3302 CDF-Chimie), puis etudie les mecanismes et cinetiques de reaction dans une gamme de temperature 100-ISO°C en mettant en evidence Ie role des impuretes sur les zones de gelification et de vitrification. Sur quelques echantillons prepares suivant des cycles thermiques difierents et en presence de difierents catalyseurs (anionique et cationique) nous avons pu etablir des relations structures chimiques proprietes thermiques.

Les memes techniques physicochimiques ont ete utilisees pour I'etude de preimpregnes (fibres de verre et fibres de carbone). Nous avons pu montrer par comparai~on avec les valeu~s obtenues sur, matrice p~re .Ie role du renfort et de j'enslmage sur les mecamsmes et cmetlques de reactlOn.

RESUL TATS ET CONCLUSION

1.- MECANISMES ET CINETIQUES DE REACTION CONSTRUCTION DU RESEAU

1.1- Mecanismes a l'etat fondu sur matrice pure

1.11- Analyse du prepolymere TGDDM

Le prepolymere tetraepoxyde commercial a ete purifie par chromato­graphie preparative (6). Le l:lroduit final et quatre impuretes ont ete isoles et analyses par RMN TH et l3C dans les Iiquides. Les resultats structuraux obtenus pour ces quatre fractions montrent que les impuretes ont des fonc­tionnalites variables et des bouts de chaInes non epoxydes qui peuvent catalyser et initier des reactions autres qu'epoxyde-amine.

1.12- Mecanisme reactionnel a j'etat fondu

Nous avons etudie la reactivite du groupement N-(epoxyde}2 a l'etat fondu vis a vis d'amines aromatiques dans une gamme de temperature 100-lSO°C sur des composes modeles (5a,0g).

Nous avons represente sur la Fig. lies difierentes reactions trouvees lors de l'etude de composes modeles et du systeme reticule.

Apres la premiere etape epoxyde-amine primaire, les reactions epoxyde­hydroxyle et epoxyde-amine secondaire peuvent etre intermoleculaires pour former des reseaux ou intramoleculaires pour former des cyCles aux chalnons non reticules.

Au tout debut de la reaction les reactions epoxyde-amine primaire sont preponderantes mais a la fin du cycle thermique quand la viscosite est grande et les diffusions limitees les reactions intermoleculaires epoxyde­hydroxyle et epoxyde-amine secondaire sont Iimitees, tandis que les reac­tions intramoleculaires de cyclisation sont favorisees.

Nous avons represente sur la Fig. 2 un suivi cinetique et mecanistique par RMN 13C solide CP MAS du systeme TGDDM-DDS apres Ie point de gel. Deux zones peuvent etre etudiees en particulier entre 140-160 pp~ {caracteristique des amines primaires et secondaires} et 30-S0 ppm {caracte­ristique des chalnons souples de la chaIne}.

37

Afin d'interpreter ces spectres nous avons synthetise des composes modEdes proches des differents chainons attend us lors de la polymerisation. Cependant etant donne les parametres specifiques de la RMN 13C solide, nous avons identifie les spectres RMN l3C liquide et enregistre les spectres solides afin de connaitre les deplacements chimiques et largeurs de raies (6). En regroupant les differentes donnees, nous avons pu selectionner les raJes caracteristiques des alcools primaires et secondaires, les ethers et les epoxydes residue Is.

A partir de ces deplacements chimiques nous avons pu deconvoluer en bandes elementaires les massifs entre 30 et SO ppm. Le carbone CH2 entre les groupements phenyles du prepolymere sert de reference interne.

Nous avons ainsi etudie differents echantillons dont les cycles thermiques ou les formulations ont ete modifiees afin de favoriser certains mecanismes comme nous l'avions montre sur un systeme type DGEBA-DDS (5H).

A partir des bilans chimiques des fonctions residue lies ou formees determines par IRFT et RMN l3C soli de (CP MAS) nous avons pu correler certaines fonctions chimiques avec les temperatures de transition vitreuse. On remarque Fi9. 3 qU'a traitement thermique et stoechiometrique identiques que les impuretes modifient la structure du reseau final (echantillons A et B). D'autre part pour les echantillons C a H, a prepolymere et stoechiometrie identiques, les modifications des traitements thermiques et I'introduction de catalyseurs anionique et cationique, modifient les mecanismes reactionnels et en consequence la structure du reseau final. Il est alors possible de correler les fonctions hydroxyle et ether formees avec les temperatures de transition vitreuse. On note ainsi que plus Ie reseau est reticule (formation des chainons hydroxyle secondaire) plus la temperature de transition vitreuse est elevee, par contre la creation de chainons ethers (reactions intramoleculaires) abaisse cette temperature.

1.21- Par microcalorimetrie Cal vet

Au point de vue cinetique no us avons suivi les reactions en isothermes par microcalorimetrie Calvet dans une gamme de temperature 100-ISO°C, Les cinetiques de reaction sui vent la meme loi cinetique que celie que nous avons proposee pour des systemes modeles c'est-a-dire autocatalyse et reaction du troisieme ordre du type dX/dt = k(R-X)(D-X)+(kIX(R-X)(D-X)ou R et D sont respectivement les concentrations en prepolymere epoxyde (R) et amine (D), pour les deux stoechiometries etudiees dans ce travail.

Si I'on porte la vitesse dX/dt en fonction du temps on remarque que dans la gamme de temperature etudiee dans ce travail les reactions d'homo­polymerisation (resine pure) sont faibles a IS0°C par rapport aux reactions mises en jeu lors de la polycondensation.

Si I'on compare main tenant la vitesse en fonction du temps d'un compose type diglycidyl aniline (DGA) polycondense avec Ie diamino diphenyl sulfone et Ie prepolymere TGDDM pur les vitesses d'initiation et d'autocatalyse sont com parables (Fig. 4a). Par contre pour Ie prepolymere commercial on remarque que les impuretes accelerent la reaction (terme d'autocatalyse) alors que les vitesses d'intiation sont com parables (Fig. 4b).

38

1.22- Par chromatographie gazeuse inverse Nous avons enregistre I'influence des impuretes sur les temps de gel et

pour cela nous avons etudie ces systemes par chromatographie gazeuse inverse. Nous avons pu ainsi montrer que la zone de gelification determinee par chromatographie gaze use inverse debute a des taux d'avancement confor­mes a la theorie de Flory (entre 25% et 32% selon la stoechiometrie) (5e,g). Par contre la gelification totale (valeur proche de celie don nee par les mesures c1assiques) peut s'observer pour des taux d'avancement tres eleves proches de 80% et depend de la purete de la resine (Fig. 5).

2.- ROLE DES FIBRES DE RENFORT SUR LA CONSTRUCTION DU RESEAU

Apres avoir identifie les proprietes de surface et les caracterisations des ensimages de differents renforts (V une etude identique sur preimpregnes (~ a celie developpee sur la resine pure a ete entreprise en presence de renforts (fibres de verre, fibres de carbone). Les -resultats obtenus au point de vue structural (construction du reseau) sont tres proches de ceux observes sur la matrice pure.

Par contre si I'on compare les cinetiques sur matrices pures et en presence de fibre, on s'aper~oit :

a) qu'elles sont, a renfort constant, sensibles au taux d'ensimage de la fibre

b) qu'en presence de renfort les cinetiques de reaction sont ralenties dans I 'ordre :

v(matrice) >v(preimpregnes fibre de verre» v(preimpregnes fibre de carbone)

Ces resultats no us semblent coherents, Ie renfort present a environ 60 % de la masse permet d'evacuer la chaleur de reaction pouvant amener une augmentation de la temperature locale. La fibre de carbone etant bien meilleure conduct rice de chaleur que la fibre de verre augmente cet effet. D'autre part, I 'ensimage plus reactif que Ie prepolymere de la resine initie et autocatalyse les reactions, modifiant ainsi les cinetiques.

References

1. J.J. King, R.N. Castonguay, J.P. Zizzi, 27th Nat. Samp. Symposium, (1982) 163

2. C.A. Cobuzzi, J.J. King, M.A. Chaudhari, 29th Nat. Samp. Symposium, (1984) 1261

3. G.L. Hagnauer,P.J. Pearce, A.C.S. Organic Coating, 46 (1982) 580 4. M.F. Grenier-Loustalot, L. Orozco, P. Grenier,

Eur. Polym. J., 22 (1986) 921 5. a) M.F. Grenier-Loustalot, F. Cazaux, J. Berecoechea, P. Grenier,

Eur. Polym. J., 20 (1984) 1137 b) M.F. Grenier-Loustalot, F. Cazaux, P. Grenier, Makromol. Chern., (1986) 1855

39

c) M.F. Grenier-Loustalot, G. Mouline, P. Grenier, O. Harran, P. Horny, Makromol. Chern. Makromol. Symp., 9 (1987) 143 d) M.F. Grenier-Loustalot, L. Orozco, P. Grenier, Eur. Polym. J., 23 (1987) 757 e) M.F. Grenier-Loustalot, G. Mouline, P. Grenier, Polymer, 28 (1987) 2275 . f) M.F. Grenier-Loustalot, L. Orozco, P. Grenier, Makromol. Chern., 188 (1987) 2559 g) M.F. Grenier-Loustalot, P. Grenier, Contrats ORET 85 34 35500 470 501

86 34 44300 470 7501 h) M.F. Grenier-Loustalot, P. Grenier, P. Horny, J. Y. Chenard, British Polym. J. (1988) (so us presse)

6. P. Horny, Contrat ORET 86 34 44300 7501

7. M.F. Grenier-Loustalot, Y. Borthomieu, P. Grenier, Surfaces and Interfaces Analysis, (1988) sous pre sse

8. M.F. Grenier-Loustalot, P. Grenier, Contrat ORET 86 34 376 000 470 7501

40

E·Ott

. I co, I coo. I f'" •

I".'~ ./c .. . ~ .c ••• ~- .. ~ .'c",_c~;c", r:: .. • I

E.SA f" . I

~ _ c •• -@- .,c ... ~_~.~ 'ata-f"-C'" ~

HUIU'III. YMLH I SA II UN ,., /c ..... --• . .. ' .... c.-, .. '.

Fig.2

" ... t ""a

Suivi cinetique par RMN 13C CPMAS

Systeme TGOOM (>96°)-005

r = I

.. LSA

150 100 50

Fig. I

Mecanisme reactionnel systeme TGOOM-OOS

18 ...

~ ____ ~------ ~l

o -so

200 Tg

180

180

170

180

150

0

R··/H

Lnl

o

200 Tg "

180 .8

180

110

.80

150

%[elher %{cHJ 20 40 80

0 20 40

Fig.3 - Correlations Tg ... (% fonctions creees)

A B DGA-DD5 A.'e/H A'B' TGDDM(96%)-DD5

A"B" TGDDM (85%)-DD5

23

A A'A" B B'B"

,(hI

Figs. 4a - b

Vitesse initiation Vitesse autocatalyse

// ~~ ./

///. / Fig. 5

o.~ /./ Zone de gelification / ./ determinee par CGI

(2) TOOO.,,151008 ,.,

25 27

41

ABSTRACT

MODIFIED BISMALEIMIDES FOR CARBON FIBRE COMPOSITES

P. KONIG, H. STENZENBERGER, M. HERZOG, W. ROMER

Technochemie Gmbh Verfahrenstechnik Gutenbergstrasse 2 - Postfach 40

6915 DOSSENHEIM - West Germany

Polybismaleimides are a class of thermosetting resins which, due to high cross-link density and aromatic backbone structure, show excellent temperature stability and outstanding hot wet performance. On account of the brittleness of the Bismaleimide resins, they have to be modified to achieve tough networks. This work describes chemical concepts to improve the fracture toughness of Bismaleimide resins, such as - addition of reactive elastomers - Michael addition chain extension - copolyaerisation with allyl terminated coaonoaers - copolyaerisation with propenyl terminated comonomers - modification with thermoplastics. Depending on the modification method, fracture toughness improvements of up to twentyfold compared with unmodified Bismaleimides can be attained.

INTRODUCTION

Carbon fibre composites are, in the meantime, being widely applied in the aircraft industry. The thermal and mechanical properties and the environmental performance of composites are, to a great extent, dependent on the properties of the resins. Epoxies, which are the most widely used composite resins, can only be used to a limited extent at higher temperatures and suffer in hot/wet conditions. It was therefore important to develop new matrix resins with improved environmental stability, temperature resistance and fracture toughness. Due to their superior temperature resistance and fire, smoke and toxicant emission properties over epoxies, Bismaleimides represent a new generation of matrix resins. Bisaaleimides, synthesized from maleic anhydride and aromatic diamines /1,2,3/, are high melting crystalline substances. Cured resin properties, like glass transition temperature (Tg), elastic modulus, moisture absorption. solubility

44

and fracture toughness are influenced by the backbone structure of the Bismaleimides and can be improved, to a certain extent, by pointed variations. Nevertheless, the Bismaleimide homopolyaers are brittle and modifications to improve the fracture toughness are required. It is a well known fact, that the refinement of one property often leads to a deterioration of other properties and consequently, the modification approach for the BMI resin must be fully understood and balanced. The following sections present five toughening concepts for Bismaleimides.

I TOUGHENING CONCEPTS FOR BISMALEIMIDES

1.1 Addition of reactive elastomers

Attempts to toughen Bismaleimides by using carboxy terminated acrylonitrile butadiene (CTBN/1300x8) rubber as a reactive elastomer compound were reported by Kinloch and Shaw 14/. The Bismaleimide resin researched was COMPIMIDE 353, which is a low melting eutectic BMI mixture. In contrast with epoxy resins, CTBN­rubber is not compatible with Bismaleimides even at high temperatures (110°C). However, during curing between 170° and 210°C the CTBN rubber reacts via double bonds with the BMI resin, producing a microphase-separated structure similar to CTBN-rubber modified epoxies. The rubber toughening increases the fracture energy significantly (Fig.l); modification of COMPIMIDE 353 with 50\ CTBN-rubber results in a fracture toughness (GIC) of 1190 J/m2. The main disadvantage of this approach is the marked decrease in the high temperature mechanical properties. The elastic modulus at 250°C shows only 0.21 GPa for the SO/50 BMI/CTBN-rubber system, which also indicates a significant reduction of the glass transition temperature. As a result of the oxidative sensitivity of the CTBN-rubber, the oxidative stability of the BMI/CTBN-rubber system is low. Nevertheless, the exceptional high fracture toughness values make this rubber­toughened BMI-system interesting for adhesive applications /5/.

1.2 Michael addition chain extension

Bismaleimides undergo a Michael addition reaction with aromatic amines 16/, amino acid hydrazide /7/, dihydrazides /8/ and many other C-H acidic compounds. Kerimid 601, a commercial product of Rhone Poulenc, is based on this concept; 4,4'-Bismaleimidodiphenyl methane (MDDM) is coreacted with 4,4'-diamino-diphenyl methane (DDM) in a non-stoichiometric ratio. A systematic study of the fracture toughness for the MDDM/DDM system /9/ indicates that the GIC reaches its maximum at a molar ratio of 111 (Fig.2). Other commercially available products, based on the Michael addition chain extension concept, are COMPIMIDE 183 and COMPIMIDE 796 from Technochemie. In this case, m-aminobenzoic acid hydrazide is used as a chain extender for the Bismaleimide. By using anionic type catalysts, like Diazabicyclooctane or Imidazole, the fracture toughness of COMPIMIDE 183 can be raised from 60 J/m2 to 120-150 J/m2. The cured Michael adducts are, in general, of lower thermal stability than the unmodified polybismaleimides. However,

45

decomposition at a significant thermal rate does not take place below 300oC.

1.3 Copolymerisation with allyl terminated comonomers

Linear chain extension is the slogan for thermosets when fracture toughness improvements are the target. An attractive toughening concept for Bismaleimides is the copolymerisation with bis(allylphenyl) compounds via a linear "ene"-type chain extension reaction, followed by a Diels-Alder addition of surplus BKI to the intermediate "ene" adduct. Couercial products are Technochemie's "toughening modifiers" COIlPIKIDE TK 120 and COMPIKIDE TK 121, as well as Ciba-Geigy's Katrimide 5292B. The "toughening modifiers" are honey-like liquids or solids at room temperature and can be easily melt-blended with 1011' melting Bismaleimides to achieve hot melt prepreg systems 110/. Their backbone structure is of high aromaticity, thus contributing to temperature resistance and high glass transition temperature (Tg). The copolymerisation of Bismaleimides with bis(allylphenyl) compounds takes place between 1200 and 2200 C. Fig.3 shows the mechanical properties of a BKI/TK 120 neat resin casting as a function of COIlPIKIDE TK 120 concentration. The fracture toughness rises up to 466 J/m2 at a toughener level of 40%, which represents ten times the value for unmodified Bismaleimides. The optimum toughener concentration for the BKI/TK 120 system appears to be between 30% and 40%. With a toughener concentration above 40' the high temperature mechanic~l properties decrease dramatically.

1.4 Copolymerisation with propenyl terminated comonomers

The bis (propenylphenoxy) compounds COKPIKIDE TK 122 and COKPIKIDE TK 123 represent two new products in Technochemie's "toughening modifiers" (Comonomers) prograue. These styrene-type derivatives copolymerise with Bismaleimides, a reaction similar to that of styrene with maleic anhydride 1111 via a double Diels­Alder addition 112/, providing tough networks. The new bis(propenylphenoxy) compounds are 1011' melting, low viscosity materials, which can easily be melt blended with bismaleimides and then cured at temperatures between 1700 and 2300 C. The correlation between the neat resin properties of a BKI/TK 123-1 system and the toughener concentration is shown in Fig.4. A low melting eutectic BKI mixture, COIlPIKIDE 796 is used as the BKI component. Flexural strength and fracture toughness increase considerably with higher comonomer concentration; the GIc-value reaches 545J/m2

at a toughener level of 30'. The glass transition temperatures (Tg) of the resin blends do not decrease significantly, in comparison to the unmodified Bismaleimide, and are between 2500

and 2600 C. At a comonomer concentration of 40% the moisture absorption decreases considerably to 2.9' vis a vis the straight BKI (4.3'). Optimum properties, considering Tg and fracture toughness are obvious between 30 and 40% bis(propenylphenoxy) compound concentration.

46

1.5 Kodification with thermoplastics

One of the most attractive ways to toughen Bismaleimides is to use engineering thermoplastics as modifiers. The idea is to combine the advantageous processing characteristics of thermosets with the high fracture toughness of thermoplastics. for this purpose two modification approaches were investigated: (a) combination of a modified Bismaleimide with a compatible thermoplastic and (b) addition of a thermoplastic with limited compatibility to a modified Bismaleimide to provide second phase toughening of the cured system. The modified BKI used was a COKPIKIDE 796/TK 123 mixture with balanced properties with respect to glass transition temperature and fracture toughness (see 1.4). It could be demonstrated, that thermoplastics, like polyetherimide (Ultem 1000) and polyhydantoin (PH-10) are excellent candidates for improving the toughness of BKI-resins /13/. Neat resin investigations of the modified resins show a 4-5 fold GIC improvement, in comparison with the COKPIKIDE 796/TK 123 resin, by addition of 20-30' thermoplastic modifier (fig.5). Vith reference to the high temperature properties, it should be noted that the Tg of BKI/thermoplastic blends is dependent on the Tg of the thermoplastic. The preferred thermoplastics have Tg's similar to the Tg of the Bismaleimide. In our series of investigations this was only achieved for the polyhydantoin. Table 1 shows composite mechanical properties of different BKI/polyhydantoin blends on T800 carbon fibres. Whereas the flexural strength decreases slightly with increasing polyhydantoin concentration, the GIC and the GIIC values increase significantly. The EDL-first-failure values increase from 126 KPa to 212 KPa for the epoxy sized fibres and from 202 KPa to 242 KPa for the polyimide sized T800 fibres, both of these representing exceptional improvements.

CONCLUSION

The five chemical concepts described to toughen Bismaleimides provide systems in which fracture toughness improvements are achieved at the expense of high temperature properties. However, modifications with Bis(allylphenyl) compounds, Bis(propenylphenoxy) compounds and thermoplastics have modification ranges, within which balanced properties are attainable. Toughening approaches could involve the combination of some of the concepts described. However the target remains the same, to achieve high toughness without any loss in temperature performance (fig.6).

REfERENCES

1. COLE, N.D. and GRUBER, V.f., US Patent 3127414, 1964 2. SEARLE, N.E., US Patent, 2444536, 1948 3. KOVACIC, P. and HEIN, R.V., J. Amer. Chem. Soc., 81

1187 (1959)

4. KINLOCH, A.J. Sci. Engng, 49

and SHAV, S.J., Amer. Chem. Soc. Polym. Kater. (1983) 307

5. SHAV, S.J. and (1985)

KINLOCH, A.J., Int. J. Adhesion Adhesives, 5(3)

47

6. BERGAIN, K., COKBET, A., GROSJEAN, P., Brit, Pat, Spec. 1, 190, 718, (1973)

7. STENZENBERGER, H.D., US Pat. 4, 211, 861 (1980) 8. STENZENBERGER, H.D., US Pat. 4, 303, 779 (1981) 4, 211, 860

(1980) 9. TUNG, C.K., LEUNG, C,L., LIAR T.T., ACS, Polym. Kat. Sci and

Engng. Vol. 52, p. 139, 1985 10. STENZENBERGER, H.D., KONIG, P., ROKER, W., HABERBOSCH, E.,

PIERCE, S., CANNING, K.S., 7th International.SAKPE Conference, Europ. Chapter Proceedings p.141 (1986)

11. BROCKNER, V., Ber. Dtsch. Chem. Ges., 75, 2043 (1942) 12. STENZENBERGER, H.D., KONIG, p" HERZOG, K., ROHER, W.,

CANNING, K., PIERCE, S., 18th International SAKPE Technical Conf., 18, 500 (1986)

13. STENZENBERGER, H.D., ROKER W., HERZOG K., KONIG P., 33rd International SAKPE Symposium 33, 1546 (1988)

ACKNOWLEDGEMENT

Part of this work was performed under contract 03M1003E5 from the West German Ministry of Research and Technology. Their support is gratefully appreciated .

... Polyhydantoln 0 0 13 13 20 33 50 Neat resin G1C (J!~) 225 225 - - 454 1090 -Fibre Size PI EP PI EP EP EP EP

90' Flex.Strength 23'C 99 92 96 91 97 91 -(MPa) 250'C 75 69 63 52 57 63 -90' Flex. Modulus 23'C 8,7 8,7 8,5 9,3 8,3 7,5 -(GPa) 250'C 7,3 9,2 7,2 7,2 6,3 6,5 -0' Short Beam Shear 23'C 103 103 101 101 96 93 -(MPa) 250'C 48 51 50 59 45 40 -0±45' Short Beam Shear 23'C 81 62 76 66 57 79 -(MPa) 250'C 43 51 50 39 44 39 -GIC-DCB (J/m2) 23'C 319 319 426 335 640 1011 1212

GUC-ENF (J/m2) 23'C - 637 824 841 705 965 1075

EDL!l:l:25)290Js 23'C

(MPa) first failure 202 126 242 132 212 176 -ultimate 571 522 569 541 517 560 -

Table 1: MechaD1cal propert1es of T800 carbon f1bre laa1Dates Resin system: COKPIMIDE 796/TM 123/PR=65/35/

xPolyhydantoin

48

1200

1000

r? 800

FM(25'C)

o

FM(250'C)

o 10 20 30 40 50

CTBN-Concenlrallon ("b.w.)

200

N~

~ ,,>1 1;;

t ~ 100

j

10 20 30 40 50

4,4'-o.amJnodiphenylmethane-Concenfration (mol %)

N~ OJ' E ~ 3

~ .c:

>0 0. co

~ I ~ -; !I ti G

~ ~

.t Ii:

0 100

0

0 10 20 30 40 50

TM 120-ConcenlraUon ("b.w.)

FS

G1C

Tg

• o

WA

Fia.1 Kechanical properties of CTBN­rubber modified bismaleimide. System:COKPIKIDE 353/CTBN

1300x8

Fia.2 Fracture energy as a function of aethylenedianiline concen­tration. System: KDDK/DDK

Fill' 3 (3 Kechanical <...

300 01 properties as a t-~

function of Bis t 2 (allylphenyl) ~

200 coapound concen-.~

t-

.~ tration. i ] System: ~ ~

= COKPIKIDE 796/

3 ~ 100 1M 120 (!5

... •

600 G,C

N' Tg .. ![ 400 ~ ::J

,j '" u.

r ~ FS 300 5i

" '" i ~ Il: 200 ~ 0 0 WA

100

10 20 30 40

1M 123-1-Concentratlon (%b.w.)

1200 O,C

1000 Tg N~ .. ~ ~ FM(25·C)

",9. 800 '" u.

i I FS 600

w

" ~ ~ FM(260·C) £ ~ 400

0 0

200

10 20 30

Thermopl881-Concenlratlon ("'b.w.)

1200 '00- - - - - - - - - - --0

90------·0

l' 1000

.3 u -0- - - - - - - - - - - - - - - - - - - -0

",- 800

i';i

~ 600 70- - - - - - - - - - -··0 '0- - - - - - ---{)

400 '0- - - - - - -40---~:: :_-_0.0

200 30------------0

L-____ --______________ --__ _ '0 10----------0

Flexural Modulus (GPa)

o Flexural Modulus and Fracture Energy al Room Temperature

o Flexural Modulus at 250'C

49

Fig.4 Mechanical

300 u .... properties as a :? function of Bis-e

4 ! i (propenylphenoxy)

I ~ benzophenone ~ 200

.§ concentration. 0-~ 1 System: COMPIMID}; ~ ,: 796/TII 123-1 3

~ ~ ...

100 •

Fia.S 300 Mechanical

e properties as a {'} function of

~ thermoplast m "- concentration. g

II ::; 200 System: COMPIIUDE u.

~ 796/TM 123/ ~ i ~ Polyhydantoin " ,:

~ ~ Ii: 100 CI ...

2 <l •

1) COMPIMIDE 796

2) MDDMIDDM (80120)

3) COMPIMIDE 796/TM 123-1 (80/20)

4) COMPIMIDE 796/TM 123 (70130)

5) BMI/TM 120 (60/40)

6) COMPIMIDE 796/TM 123-1 (70/30)

7) 8MI/O,O' ~Dianylblsphellol A (60/40)

H) COMPIMIOE 7961lM 1231

Polyelherllnido (48,2/25,9/2t>,9)

9) COMPIMIDE 796ITM 1231

Polyhydantoin (37120/33)

10) COMPIMIDE 353/CTBN 1300'8(50/50)

Fig.6: Neat resin fracture energy (GIC) versus flexural modulus

ENHANCED BONDING OF FIBER REINFORCEMENTS TO THERMOSET RESINS

G. SUGERMAN, S.M. GABAVSON', W.E. CHITWOOD', S.J. MONTE

Kenrich Petrochemicals, Inc. 140 East 22nd Street, 07002-0032 BA YONNE New Jersey - USA

"General Dynamics PO Box 85990 - CA 92138-5990 SAN DIEGO - USA

Abstract

Improved bonding of fiberglass. Kevlar and carbon fiber to a variety of thermoset resins has been achieved via the addition of minor amounts of organotitanate and/or organozirconate coupling agents. Peel strength. fiber pullout and chemical resistance data is provided to demonstrate significant improvements in bonding between these fibers and epoxy. polyurethane. polyester and vinyl ester resins.

1. INTRODUCTION Most fiber reinforced composites fail due to inadequate bonding

at the interface between reinforcement and matrix resin. The causes for failure are often a combination of system inhomogeneity consequent to processing limitations. poor wettability. and a lack of chemical compatibility/reactivity between the fiber and matrix. Organosilanes have long been utilized to enhance the chemical bonding of thermoset resins with silaceous (glass) surfaces; however. organosilanes are essentially non-functional as bonding agents for graphitic. organic and/ or metallic fibers. Furthermore. silanes usually confer minimal processability benefits-even in fiberglass reinforced formulations. thereby requiring the incorporation of diluents. and/or process aids which often diminish performance.

Enhanced processability and improved bonding to silacious and non-silacious reinforcements may be achieved by the use of organometallic titanium and zirconium coupling agents. When added at very low levels. 0.1 to 0.5 wt. percent on formulation solids. they typically improve substantially processability and dispersion as a consequence of molecular wetting and upgraded matrix to fiber bonding as a result of specific chemical interaction with each component at the interface. This paper demonstrates that the use of organometall ic

52

coupling agents. optionally in conjunction with organosilanes for fiberglass reinforced composites only. produces a wide variety of fiber reinforced composites with superior processability. physical properties and environmental stability.

2. DISCUSSION AND RESULTS Performance evaluations of organosilane. titanate and zirconate

coupling agents (cf. Table 1) as l~ level additives were made in: aramid and carbon fiber reinforced epoxy. polyurethane and vinyl ester. and in fiberglass reinforced epoxy and polyester resins. Pull out energies (adhesion) for long fiber reinforcements in short fiber modified matrix resin. and processability and physical properties of the modified matrix were determined independently. Results of original and ten day 10% salt water boil tests are detailed in Tables 2 through 5.

It was found that the introduction of appropriate organotitanate and/or organozirconate coupling agent(s) at levels of 0.2 to 0.5 parts per 100 parts of resin substantially improved processing. original and aged physical properties in each system tested. Only in fiberglass reinforced formulations were organosilanes found efficacious; even in these systems. they provided little. if any. rheological benefits. and were found to produce physical property enhancements and corrosion protection inferior to those available from their use in combination with amino zirconates. In epoxy. polyester. vinyl ester and polyurethane resin systems. amino zirconates appeared to provide maximal benefits. in terms of physical property enhancements and corrosion protection (50 to 100% in original properties and an order of magnitude corrosion loss minimization enhancement after 10 days of 10% salt spray at lOOOe). Their use in conjunction with coordinate type phosphite adduct analogs provided greater enhancement of processability. and in many instances. further improved original property retention under severe salt spray conditioning. Additional data relating to the usage of organometallic coupling agents is summarized in a comprehensive Reference Manual (1).

3. EXPERIMENTAL A fiber pullout evaluation of 80 cm long. 2 mm (nominal)

diameter twisted multistrand fiber were clamped vertically under 10 joules of tension under a 20 cm long by 2.5 cm diameter vertically split and clamped teflon cylinder centered axially on the fiber. The cylinder was filled with the short fiber reinforced resin composite. optionally containing the additive which when present was introduced in advance of fiber. The tightly filled cylinder was then cured as specified. the Teflon hemicylinders removed. and the cylinder sawed in half across its short axis. Each resulting half cylinder and protruding fiber were separately clamped and separated by axial stretching until fiber separation from the resin matrix was achieved. Maximum yield stress was recorded for 10 samples of each type and values within + 5% of the arithmetic mean averaged.

In no-case were data from more than one sample discarded as a consequence of this limitation. Physical evaluations of the short fiber reinforced base resin compositions were made separately from 15 cm x 1 cm x .3 cm dogbone castings cured as specified in the pullout

53

experiments. equilibration to boiling 10

All samples were tested after 24 hours of ambient both prior to initial testing and subsequent to exposure percent saline solution in the aging experiments.

4. CONCLUSION The data presented in Tables 2 through 5 demonstrate clearly

that the processability, and both initial and aged physical properties of fiber reinforced thermoset resins may be substantially enhanced by addition of low levels of appropriate organometallic coupling agents. In thermoset resin applications involving carbon fiber, fiberglass and aramid reinforcement, aminofunctional zirconates, optionally in combination with organofunctional silanes as coagents or with coordinate type organometallics as rheology enhancers, appear to provide substantial adhesion, processing and anticorrosive benefits.

It would appear that the usage of organotitanate and/or organozirconate coupling agents, either independently or optionally in conjunction with other surface reactive materials, such as organosilanes in the speci fic case of fiberglass reinforced systems, provides an inexpensive, highly functional methodology for drastically improving the processability, physical properties and corrosion resistance of fiber reinforced thermoset resin composites.

5. REFERENCES (1) Monte, S. J. and G. Sugerman, Ph.D., Kenrich

Petrochemicals, Inc., "Ken-React® Reference Manual - Titanate, Zirconate and Aluminate Coupling Agents," (Bulletin KR 1084L-2), 230 pages, (February 1987 - Revised Edition).

Code A-174 A-187 A-llOO KR 55 adduct 2

KR 134S diolato

Nomenclature

TABLE 1 NOMENCLATURE

3-Methacryloxypropyl, trimethoxy silane 3-Glycidoxypropyl, trimethoxy silane 3-Aminopropyl, trimethoxy silane Titanium IV tetrakis(bis 2-propenolato methyl)-I-butanolato

moles (di-tridecyl)hydrogen phosphite Titanium IV bis[4-(2-phenyl)propyl-2]phenolato,oxoethylene-

KZ 55 Zirconium IV tetrakis (2,2-bis propenolato methyl)butanolato, adduct with 2 moles bis tridecyl

hydrogen phosphite LlCA 44 Titanium IV neoalkenolato, tris (2-ethylenediamino)ethylato LZ 37 Zirconium IV (2,2-bis propenolatomethyl)butanolato, tris 4-amino-benzoato-O LZ 44 Zirconium IV neoalkenolato, tris (2-ethylenediamino) ethylato LZ 97 Zirconium IV neoalkenolato, tris (3-amino)phenylato

54

TABLE 2 EVALUATION OF VARIOUS COUPLING AGENTS IN FIBER REINFORCED (ORIGINAL/AGED PROPERTIES) EPOXY COMPOSITES

Formulation Parts By Weight Resin, DEN 438 (Novalak - Dow) 100 Hardener, Methyl Nadic Anhydride 87.5 Fiber (Short - 5 mm) 40 Additive(s) as shown Brookfield Viscosity measured @ 80°C; Cure 30 min. @ 150°C. Postcure: 4 hr. @ 180°C. Aging 240 hr. in boiling 10% aqueous salt solution. A) Carbon Fiber Reinforcement (IM-6, Hercules)

Lonfj Fiber Short Fiber Vehicle Only Tensile Flexural Compressive Falling Ball

Additive: Pullout Strength Strength Impact Str. (PEW) Energr J fjPa fjPa kJ/m

Control 62/21 0.94/0.41 1. 72/1.07 1. 5/0. 5 LZ 97:0.4 119/113 1.24/1.13 2.71/2.49 2.3/1.8 LZ 37:0.3 122/119 1.42/1.34 2.96/2.70 2.6/2.4 KR 55 :0.2 78/62 1.31/1.20 2.07/2.01 3.1/2.7 KR 55:0.3 97/81 0.24:0.17 2.28/2.15 3.4/2.7 KR 55:0.3 + LZ 37:0.3 151/148 1. 72/1. 67 2.82/2.69 2.9/2.8 KR 55:0.'1 + LZ 37:0.4 132/130 1. 59/1. 46 2.61/2.48 2.6/2.5 B) Fiberglass Reinforcement (E-Glass, Certainteed)

Lona Fiber Short Fiber Vehicle Qnll Tensile Flexural Compressive Falling Ball

Additive: Pullout Strength Strength Impact Str. (PEW) Eneral J aPa gPa kJ/m

Control 54/14 0.82/0.40 1.54/1.21 1.7/1.0 A-1100:0.4 63/41 0.96/0.69 1.69/1.43 1.9/1.5 A-187:0.4 65/37 0.92/0.74 1.58/1.51 1.9/1.6 LZ 37:0.4 113/81 1. 27/1. 28 2.17/2.09 1. 8/1. 8 KR 134S:0.4 79/76 0.96/0.92 1.81/1.77 2.5/2.4 LZ 37:0.2 + KZ 55:0.2 117/113 1.42/1.36 2.16/2.09 2.8/2.4 LZ 37:0.2 + KR 134S:0.2 126/123 1.37/1.31 2.23/2.11 2.9/2.7 LZ 37 :0.2 + A-187:0.2 129/121 1.61/1.50 2.32/2.23 2.1/2.0

C) Aramid Reinforcement (Kevlar, DUEont) Lona Fiber Short Fiber Vehicle OnlX

Tensile Flexural Compressive Falling Ball Additive: Pullout Strength Strength Impact Str. 0.4 (PEW) Eneral J fjPa sPa kJ/m Control 59/38 1.16/0.69 2.32/2.03 1. 9/1.1 A-1100 62/49 1.29/1.14 2.42/2.23 2.1/1.3 LZ 97 76/67 1. 41/1. 27 2.84/2.58 3.5/2.8 LZ 37 84/72 1. 42/1. 38 2.79/2.67 3.2/2.7 KR 134S 73/67 1.31/1. 27 2.27/2.14 2.9/2.9

55

TABLE 3 EVALUATION or VARIOUS COUPLING AGENTS IN rIBERGLASS (ORIGINAL/AGED

PROPERTIES) REINFORCED POLYESTER Formulation Parts By Weight Resin (Stypo1 40-1029 - Freeman Chemical) Curatives: Benzoyl peroxide (Certainteed) bis t-buty1peroxycyc1ohexane Short (5mm) and long fiber (fiberglass) Additive(s) Brookfield viscosity measured @ 25°C. Cure 20 by postcure of 2 hr. @ 180°C.

min.

100 0.5 0.5 40

as shown @ 150°C followed

Long Fiber Tensile

Additive: Pullout

Short Fiber Vehicle Only

(PBW) Energy J 42/27 49/41 50/44 59/56 62/58

Control A-1100:0.4 A-174:0.4 LZ 97 :0.4 LZ 37:0.4 LZ 37 :0.2 +

Flexural Modulus gPa

16/11 19/14 18/16 21/17 23/21

Flexural HDT Strength 1.18 mPa gPa °C

0.83/0.49 201/169 1.05/0.88 218/194 1.06/1.01 217/203 1.14/1.12 221/216 1.27/1.19 229/216

N.l. kJ/m

1.15/0.89 1.32/0.97 1.28/1.10 1.56/1.47 1. 23/1.09

KZ TPP:0.2 59/55 LZ 37:0.2 + A-174:0.2 71/69

24/23 1.34/1.27 242/228 1.29/1.17

28/26 1.46/1.42 279/257 1.42/1.38

TABLE 4 EVALUATION or VARIOUS COUPLING AGDTS IN rIBER REINFORCED VINYL

ESTER (ORIGINAL! AGED PROPERTIES)

Formulation Parts By Weight Resin. Vinyl Ester (Dion VER 9400 - Koppers) Bispheno1 A dimethacry1ate (SR 348 - Sartomer) Catalyst. benzoyl peroxide

bis t-butylperoxycyclohexane Short (5 mm) fiber Long Fiber (Carbon - IM-7 - Hercules) Additive(s). as shown

80 20 0.5 0.5 40 40 0.4

Brookfield Viscosity measured @ 50°C. Cure 20 min. @ 150°C followed by postcure 4 hr. @ 210°C. Aging 240 hr. in boiling 10% aqueous saline solution. A) Vinyl Ester/Carbon Fiber

Long Fiber Pullout

Additive Energy 0.4 (PBW) --=-:.::(J~)~ Control 56:24 A-1100 54:31 A-174 58:30 LZ 97:0.2 + A-174:0.2 LZ 37:0.2 + A-174:0.2

70:67

70:61

Short Fiber Vehicle Only Flexural Flexural HDT

Modulus Strength 1.18 mPa gPa gPa °C 18/12 1.91/1.13 208/170 19/13 1.90/1.05 207/167 19/11 1.96/1.09 210/173

N.!. kJ/m 1.37/0.90 1.32/0.94 1.39/0.97

21/20 2.18/2.07 234/230 2.19/2.08

20/20 1.14/1.08 231/227 2.01/1. 92

56

B) VinI1 Ester Aramid Long Fiber Short Fiber Vehicle On1l

Pullout Flexural Flexural HDT Additive Yield Modulus Strength 1.18 mPa N.I. 0.4 (PBW) (kg) sPa SPa °C kJ/m Control 53:21 20/12 1.82/1.07 211/162 1.42/0.85 A-1100 57 :32 20/14 1.85/1.26 219/174 1.54/0.89 A-174 55:24 19/11 1. 91/1.14 208/177 1. 47/0.90 LZ 97 86:73 24/21 2.86/2.42 236/221 2.31/1.96 LZ 37 97:86 28/25 3.47/2.89 228/216 2.42/2.31 LZ 97:0.2 + A-174:0.2 84:70 21/17 2.69/2.31 214/201 2.09/1.93 LZ 37 :0.2 + A-174:0.2 92:85 25/23 3.05/2.63 209/197 2.16/1.95 LZ 37:0.2 + KZ TPP:0.2 116/109 28/27 3.78/3.46 231/219 3.02/2.87

TABLE 5 EVALUATION OF VARIOUS COUPLING AGENTS IN FIBER (ORIGINAL/AGED

PROPERTIES) HEINFriRCED POLYURETHANE Formulation Parts By Weight Resin. Po1yo1 (Poly BD-45HT - Sartomer) 85

(Isono1-100 - Dow) 15 Isocyaante (PAPI-94 ..; Dow) 28

Catalyst (T12 - Air Products) 0.5 Short Fiber Reinforcement (as shown) 40 A) Carbon Fiber (IM-7 - Hercules)

Long Fiber Tensile

Additive: Pullout (PBW)

Control A-1100:0.4 A-187:0.4 LZ 97 :0.4 LZ 37:0.4 LICA 44:0.4 LZ 37:0.2 + A-187:0.2

Enern J 32/21 43/28 37/26 52/47 59/49 52/37

59/54

Flexural Strength

gPa 1.02/0.63 1.05/0.79 1.05/0.86 2.37:1.82 1. 91/1. 63 1.97/1.42

2.19/2.08

B) Aramid Fiber (Kev1ar - DuPont) Lons Fiber

Tensile Flexural Additive: Pullout Strength

(PBW) Enern J gPa Control 39:24 1.21/0.72 A-1100:0.4 46:31 1.24/0.83 A-187:0.4 41:29 1.19/0.80 LZ 97:0.4 57:51 1.63/1.51 LZ 37:0.2 + KZ 55:0.2 62:57 1.97/1. 82 LZ 97 :0.2 + KZ 55:0.2 60:58 1. 88/1. 79

Short Fiber Vehicle Onll Compressive Falling Ball Strength Impact Str. sPa kJ/m 1.32/0.74 3.9/2.4 1.41/0.82 4.0/2.9 1.53/0.89 4.1/3.3 2.07/1.89 4.8/4.6 2.16/2.01 5.1/4.7 1.94/1.57 4.7/3.6

2.41/2.29 5.4/5.0

Short Fiber Vehicle Onll Compressive Falling Ball Strength Impact Str.

gPa kJ/m 1.56/0.91 4.3/2.1 1.71/0.96 4.4/2.4 1.67/0.97 4.4/2.6 2.61/2.39 5.4/5.1

2.86/2.71 5.4/4.9

2.67/2.53 6.2/5.8

INFLUENCE OF THE THICKENING AGENTS AND SOME EXTERNAL PARAMETERS TO

THE FORMULATION ON THE VISCOSIMETRIC KINETICS OF THE PREIMPREGNATED POLYESTER

ABSTRACT

A. VALEA PEREZ, M.L. GONZALEZ, I. MONDRAGON'

Escuela Universitaria De Ingenieria Tecnica Industrial Plaza de la Casilla 3 - 48012 BILBAO - Spain

'Escuela Universitaria de Ingeniera Tecnica Industrial Plaza de Pio XII - SAN SEBASTIAN - Spain

The influence of such thickening agents as CaO; MgO; Ca(OH~ and Mg(OH) on the viscosimetric behaviour of preimpregna­ted polyester samp~es (SMC), has been studied. The results of the­se studies, allow us to conclude that, even though with all the agents the viscosity increases, the MgO is the most effective, as it originates the faster kinetics of viscosimetric increase, com­patible with most of the references of transforming through usual moulding. In addition we try to present and discuss some results we have obtained to prove the influence of amount of water and other of the pre impregnated SMC preparation.

INTRODUCTION In the last years, the classic materials such as

rolled steel, aluminium, wood and glass have been progressive and increasingly substituted by plastic, reinforced or not, as a con­sequence of their excellent service/cost ratio. At the same time, the enormous possibilities when designing, makes these plastics compulsory in nowadays engineering.

Wi th respect to the reinforced plastic group, the so called nonsaturated polyester composites, reinforced or not with fiber­glass, have been widely used since 1955 due to their applicability for the manufacturing of large series.

In this paper we will study the aging or thickening agents, whose aim is to give to the pre impregnated the minimum necessary consistency to allow the moulding by compression (1-3). In fact, we will study the influence of some of the most important ones on the viscosimetric behaviour (4) of pre impregnated polyester (SMC). The importance of these studies is based on the produce an optimal cycle.

58

EXPERIMENTAL

We have chosen a formulation that, at present, is used at technical level (5), but not using addi ti ves that could produce marginal reactions.

Likewise, we have not used reinforcement of fiberglass, becau­se besides that it is not necessary for our aims, it.would greatly raise difficulties for the realization of the viscosimetric tests as well as for its interpretation. In resume, our procedure is ba­sed in a polyester resin (BASF), constituted by a bicomponent sys­tem 60/40 of Palatal KR5510/5511 -(PS), whose initial viscosity is of 3500 cps under test conditions, a neutral load constituted by calcium carbonate (OMYA), and thickening agent.

The preparation of the pre impregnated sample was made mixing the three components and after the mixture was strongly agitated through a Raynerie turbin (at 1200 rpm) for 5 minutes. Finally, the mixture was thermostated at 293 +0,1 K.

The viscosity determinations were made using a viscosimeter Brookfield model RVT with 8 different speeds. The highest viscosi­ty values were obtained through the consistency, measuring the breakthrough during 5 seconds of the calibrated top of a penetro­meter Vaschetti (ASTM D-217).

The'following thickening agents were used: calcium oxide, mag­nesium oxide, calcium hydroxide and magnesium hydroxide. The used ratios were in all cases 1,5% (w/w) in relation to the resin.

The whole amount of water was measured by Karl-Fischer titra­tion in a Karl-Fischer apparatus of Crison (ASTM method). The con­trol of temperature was made by digital Crison thermometer 616/N with a platinum sonde, which was previously calibrated in the ran­ge 0-160 o +0,05°C.

In the study we have taken a mixture system of 0,21% w. of to­tal water (it is named Solution I) and adding water, we have sys­tematically prepared these next solutions: Solution II (0,43% w.of water); Solution III (0,68% w H20) and Solution IV (0,93% w H20), which have been measured by Karl-Fischer titration. So the diffe­rences of these solutions are the weight of water. Analogous pre­cautions have been taken into the study of the influence the other parameters, so, formulation and morphology are also identical in all cases.

RESULTS AND DISCUSSION

In figure 1 the kinetics representation of the viscosimeter increase is shown. In the case of the CaO and MgO, the values from the discontinuity are given in Vaschetti pe­netration units ( 1 unity = 1/10 mm). As it can be seen, the gra­phics are the typical (4) for a preimpregnated formulation SMC, as it is our case. We can clearly distinguish three different zo­nes. After analysing Figure 1, we can deduce that the magilesi urn oxide allow us to obtain a faster thickening kinetics, followed by the calcium hydroxide. However, after 21 days (storage phase)

59

this behaviour changes, as the viscometric increase is higher with CaO than with Ca(OH)2'

We have to emphasize the high viscosimetric increase produced between the second and third day in the case of preimpregnated with CaO, reaching thickening speeds of 7273 poisses/hour, that makes cri tical the control. The increase of the viscosity obtained in all cases is a consequence of the reaction between the polyester final groups and the thickening agents according to the scheme (5, 8) in figure 2.

Table 1 shows the characteristical indexes of every phase (impregnation, aging and storage) for the pre impregnated obtained using differents thickening agents. After studying these values, we deduce that there is a drastical difference between the final preimpregnated products (periods higher than 22 days). This diffe­rence will oblige us to keep a careful control of the transforma -tion as the mouldabili ty and plasticity characteristics are very different from one case to another, depending upon the thickener used in the preparation.

In relation to the viscometric increase we can say that the best results are attainable with MgO in such a way that being com­patible to a good impregnability we obtain a shorter machine pe­riod which means an important saving.

The influence that the amount of water coming from loads, re­sin, and thickening may exert on viscosimetric growth kinetic of preimpregnated SMC has been studied because it has been experimen­tally confirmed that changes in moiety play an important role on kinetic ageing of preimpregnated SMC, arising some fundamental di­fferent behaviours in the last phase of its transformation.

In figure 3 are summarized the viscosimetric results obtained for four different formulations, with the same thickening agent MgO (at 1,5% w. on resin) but different amounts of total water above cited. Anyone can see that the kinetic behaviour seems to be so complicated. So as to find a law of variation we have parti­ally considered every typical phase that a pre impregnated has to suffer, they were mentioned in the previous work. In table 2 are summarized the experimental values obtained for the characteristic index of every phase. Looking at these values, one can clearly no­tice three things: 1) Only IV 100 seem to arise wi th the whole amount of water. 2) The absolute differences of viscosity seem to shorten in a certain time, showing that the IV 21 index is similar for all the pre impregnated chosen which are being investigated. 3) It seem not to be a systematic viscosimetric behaviour of the pre impregnated chosen during the impregnation phase or it would be better to say that in the taken range of total water there is

a, machine-cycle, minimum time to the 12 index which will be in­teresting to produce SMC. Obviously, this conclusion has to be confimed in more experimental works to get an adequate recurrence equation considering a large number of dates.

To appreciate fairly the variation of the different characte­ristic index with the whole amount of water we have investigated

60

an easiest equation between 12 index (minutes) and the amount of water (x=% w, on resin), using a computer, it is:

2 12 (minutes) = 162,642·x - 169,825'x + 107,519

with correlation coeff.=0,99993 and determination coeff.= 0,99986. This behaviour might indicate that during the pre impregnation

phase and until it had got a certain level of water, it performs the same role as a catalyst of reaction.

Exactly the same came be argued about the variation of IV21 in­dex, which characterizes the storage phase; in this case it also seems to be a concentration at about 0,6-0,7% w. in water, where the product has a minimum penetrability. If we take together both characteristics we can deduce that increases of the total water in about 0,55-0,70% w. in our system will be related to minimum production time-machine and the final preimpregnated SMC will have a minimum penetrability (that is, rigid products and less handling during moulding) after 21 days.

About the variation of the IV21 index with the amount of water the following equation have been found as the simplest:

IV21 (units) = 13,3498'x2 - 16,0109'x + 43,8121

with correlation coeff.=0,9954 and determination coeff.=0,9909. So during the ageing phase, if the whole amount of water in­

creases, the reactivity will disminish. The variation of IV100 number with the total amount of water,

we can deduce that there are a lineal and simple relation between IV100 index (hr) and the amount of water. Taking an adjustment by computer, we have attained the next regression equation:

IV100 (hr) = 60,9831·x + 18,6520

with correlation coeff.=0,99994 and determination coeff.=0,99988. From the great slope of this line can be assured that the

amount of water plays a fundamental role on the required ageing time that the pre impregnated SMC needs to obtain tha adecuated vis­cosity so as to be moulding in usually conditions.

It also seems to be clear that you have to know the total amount of water into a specific formulation in order to obtain the same characteristic and reproducibility in the viscosimetric growth kinetic of different lots of fabrication.

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FIBRES

Chairman: Dr A.R. BUNSELL Ecole Nationale Superieure des Mines de Paris

ABSTRACT

THE MODULUS OF ALUMINA FIBRES CONTAINING MESOPORESDEPENDENCE

ON ORIENTATION DISTRIBUTION

M. H.STACEY

ICI Advanced Materials The Heath - WA74QE RUNCORN. Cheshire - England

Sol-gel ~-alumina fibres containing a proportion of axially-oriented pores have been compared with others containing only random porosity. The former fibres possess better mechanical properties. A simple theory is presented to account qualitatively for these effects which are related to the optical birefringence of fibres containing some axially oriented pores.

INTRODUCTION

Alumina fibres made by sol-gel routes invariably contain pores after decomposition. A study of the sintering of such pores in experimental 10 ~ diameter fibres concluded that the pores initially form a random network of cylinders. Upon sintering it seems that axially-aligned pores may be preferentially eliminated since the modulus increases at a lower rate with respect to porosity than was reported for macroscopic bars of sintered a-alumina powder (1) .

This paper addresses the questions of how the mechanical properties of porous alumina fibres made by a sol-gel method are influenced when the pores are organised in other ways; and specifically when some pores are aligned parallel to the fibre axis. In order to do this fibres having a range of diameters were generated containing wholly random pores (Samples A and B) and

66

alternatively fibres were made in which a fraction of the pores were aligned with the fibre axis (Samples C and D). Each type of fibre is fully characterised in respect of its' porosity and mechanical properties.

EXPERIMENTAL

Batches of aligned multifilament alumina fibre tows were made by aproprietory version of the sol-gel process in which the conversion to ceramic fibre was halted when the fibres were found by powder X-ray diffraction to be polycrystalline ~-alumina. Usually such fibres contain about 40% total porosity.

Porosity was characterised by nitrogen adsorption isotherms at 77K, by Small-Angle Neutron Scattering (SANS) (2), and by measurement of optical double refraction. Fibres were ground finely, and degassed at 250·C at 10- 3 mbar before loading into a Digisorb nitrogen adsorption apparatus. For neutron scattering the fibre tows were loaded as made into IOmm i.d. fused silica tubes and degassed as for the nitrogen adsorption, the tubes being sealed under vacuum. Ca 0.2g/cm fibres were loaded which resulted in transmission coefficients in the range 0.6-0.B (2).

For optical measurements a Nikon polarising microscope was fitted with a Senarmont compensator and a few fibres were mounted dry under a cover slip. Fibres were usually observed at x200 or x400 magnification. Some of the batches of fibres were observed to be positively uniaxial double-refracting (i.e. fibre axis and optic axis are parallel) and the compensator allowed the measurement of the path difference for individual fibres. The refractive index difference was calculated for observations on ca 20 fibres from the formula

SA

IBO.diameter

where S = angle of rotation of the analyser which caused maximum darkening of the fibre image and A=550 nm.

The mechanical properties of single fibres were measured using a Marsh microtensile machine (3). The fibres (lor 2mm gauge length) were selected by hand and glued to the anvils with molten diphenylcarbazide. Usually about 20 fibres were tested from each sample. The fibre diameter was measured after the fibre had broken by transferring the broken end(s) to a slide and mounting on a Nikon polarising microsope fitted with an eyepiece vernier graticule using a magnification of XIOOO. Sizing errors are ca 0.2~. The strength and modulus of each fibre was calculated from the force/extension graph.

67

RESULTS

POROSITY CHARACTERISATION

Some batches (C and D) of porous alumina fibres were observed to be strongly double-refracting in polarised light, whereas other batches (A and B) were optically inactive. This was determined to be a case of form birefringence since the ~-alumina phase is essentially cubic in symmetry and so cannot be the source of double-refraction. Observations showed that each fibre was uniformly double-refracting over many mms length and that for any given batch the path difference was roughly independent of diameter. It follows from equation (1) above that the refractive index difference is therefore inversely proportional to fibre diameter. Typical results are shown in Fig 1 for the refractive index difference as a function of fibre diameter for sample C. It was also shown that if the fibres were immersed in an oil of refractive index 1.6 then the double refraction was eliminated. It follows from Wiener's theory of form birefringence that fibrillar pores parallel to the fibre axes are responsible (4).

Independent measurements of the total fibre porosity were determined from the nitrogen adsorption isotherms at 77K (fig 2). The isotherms are typical of mesoporous materials having type H2 hysteresis loops. They exhibit a well-defined maximum pore volume above relative pressures of ca 0.8. Pore diameter distributions can be calculated from the adsorption loops using BJH theory in which cylindrical pores are assumed. Mean values are given in Table 1 as the distributions were very similar in all cases. Fig 3 shows that the birefringent fibres also had anisometric SANS patterns in which scattering perpendicular to the fibre axis was 2-3x more intense than that parallel to the fibre axis. On the other hand inactive fibres also had isometric SANS patterns. In general those batches of fibres containing only random pores had higher surface areas and porosities than those exhibiting double refraction but mean pore diameters were ca 5 nm in all cases. (see Table 1).

Thinned axial sections of fibres were prepared by Ar atom etching of bundles embedded in epoxy resin. These were examined in a Phillips EM400 at lOOK magnification. The fibre batches exhibiting double-refraction had a marked linear texture consisting of bright features having a width of ca Snm parallel to the fibre axis being spaced at 10 nm intervals (see Fig 4); some randomly orientated texture due to mesoporosity was also visible. Optically inactive fibres only possessed the random texture (Fig 5).

The essential difference between the two classes of materials is that when optically active the mesoporosity is in part aligned with the fibre axes. The optical birefringence provides an empirical measure of the fraction of pores which are thus aligned. The SANS observations are also in accord with this interpretation but a full quantitative analysis of the SANS results is not practical at this time and will be dealt with in a future paper.

68

2. MECHANICAL PROPERTIES

Two batches of random pore fibres (samples A and B) had an insignificant variation of mean modulus with fibre diameter. Fig 6 shows the scatter plot of strength and modulus for fibres having a total porosity of 40% and with fibre diameters in the range of 3-12 ~. The data show a decrease in strength with fibre diameter which when fitted to the Weibull equation

Failure Probability = exp (-Volume am) gave a value for m of 7.5+-0.8.

Strongly birefringent fibre batches (Samples C and D) gave the results in Fig 7. As can be seen there is a strong variation in results with fibre diameter which parallels the variation in double refraction between fibres with diameter. At large diameters the results converge with the results previously found for random pore batches while at low diameters a high maximum value seems to be found. Though it is possible to fit empirical regression equations to such data this is not very profitable since a single data set does not cover a wide enough diameter range. Attempts to fit a modified Weibull equation to the strength data by fitting a diameter exponent other than 2 were not very successful since the required assymptotic behaviour cannot be simulated.

DISCUSSION

Previous research has shown that when large diameter random pore aluminas are sintered the modulus and strength do not increase as sharply with decreasing porosity as do sintered powder bars of a-alumina. I have proposed that in sol-gel aluminas axially-aligned pores sinter slightly more rapidly than radially-aligned pores to account for this (1). In this new work, the existence of preferred axial orientation mesopores in some batches of fibres raises new questions. Specifically how can the modulus and strength of such fibres be increased so much relative to totally-random pore batches. Also whether the sintering behaviour of such aligned pore fibres differs or not will be reported later.

It is clear that a quantitative microstructural model of fibres is needed to aid discussion. The data obtained from the analytical techniques clearly show that birefringent fibres contain both random pores and a fraction of axially-aligned pores all of which are in the 2-10nm diameter range. From Wiener's theory of form birefringence (4), porous ~-alumina having a totally fibrillar porosity of 30-40% should have a maximum refractive index difference of ca 0.04. The range of values found (Fig 1),0.01-0.03, indicates that 30-80% of the porosity in sample C is fibrillar depending on the fibre diameter.

Wang (5) has shown that the modulus of sintered a-Alumina

69

powder decreases exponentially with porosity whereas if porosity is present as axially-oriented fibrils then the rule of mixtures should apply (6). Although samples C&D are less porous than samples A&B the increase in modulus and strength is much greater than would be expected if their porosity is totally random. Since the modulus decrease expected from the rule of mixtures is much less than that given by Wang, I conclude that the presence of some axially-oriented pores in samples C&D is the reason for their improved mechanical properties.

A mathematical model will be reported in a later paper which enables quantitative predictions of birefringence and modulus for ~-Alumina to be made for fibres having pores which are partly random and partly axially orientated.

REFERENCES

1 - Stacey M H., Science of Ceramics 14 (1987) 291-297 2 - Stacey M H., in "Characterisation of Porous Solids"

(Unger K.K. et al. Edrs) 55-65 1988, Elsevier, Amsterdam 3 - Marsh, D M., J. Sci. Inst. 38 (1961) 229-234 4 - Partington JR." "An Advanced Treatise on Physical

Chemistry" IV, p275, Longmans, London, 1953. 5 - Wang J C., J. Mater. Sci. 19 (1984) 809-814. 6 - Kelly A. and Macmillan N H. "Strong Solids" 3rd Ed (1986).

p242, Clarendon Press, Oxford.

TABLE 1 Mean Properties of Samples

Sample Diameter Gauge Total Pore True Double Modulus Strength length por. dia. Density Ref.

)JIll mm % nm g/ml GPa MPa

Random pore A 4.76+-0.8 1 41.8 6.2 3.23 nil 58+-9 612+-121 B 8.64+-1.8 2 38.8 5.6 3.17 nil 54+-12 431+-136

Aligned pore C 2.62+-0.4 31.2 5.4 3.17 0.0216 151+-50 1486+-600 D 8.26+-1. 1 22.3 4.8 3.08 +(na*) 102+-27 910+-210

Footnote: "na=sample positively double-refracting but not quantified

70

<II <.l C <II ,.. <II .....

..... 'rl

"" x <II

"" c .... <II :> .... ... <.l <II ,..

..... <II

<>:

""' ., ...

0.03

0.02

0 . 01

Fig

glOOO <.l ...... :>.. ... .... ~ 100 <II u C

H

10

. !

• "

2 3 diameter/jJm

Sample C Refractive index differences

log (Q/K):4~Sine/A Fig 3 SANS for samples A and C

~1000 ...... .. : .. b a ...... ,., •

.. • •• tI' . .. . . (1) 400 • .. " 0 .... .... . •

200 <11

Po. 100 C-' ...... , ILl . : " .. ......... . ........ . (1) • j. . 0 40 ....

2 4 6 8 10 diameter/jJm

Fig 6 Strength and Modulus • o. Sample A .. I:> Sample B

Po. E-< OIl

u <11

N Z

.... :>::

o 0.2 0 . 4 0 . 6 PIPo

0.1"-

0 . 8 1.0

! Fig 2 Nitrogen isotherms at 77K

Fig 4 TEM Sample C Fibre Axis Vertical

Fig 5 TEM Sample A

3000 ~~ . •

1000 DO," 0

" °Drllo t,,"~

" 500

, '. :, .. : .. ~ .... 100 . ' ... : . .. .. 50

2 4 6 8 diameter/jJm

Fig 7 Strength and Modulus • o . Sample C .. I:> Sample

"

10

D

STATISTICAL MECHANICAL BREAKDOWN OF SINGLE FIBRES AND MICROCOMPOSITES USING VIDEO

MICROPHOTOGRAPHIC TECHNIQUES

ABSTRACT

H. D. WAGNER, L. WIM STEENBAKKERS'

The Weizmann Institute of Science Department of Materials Research

76100 REHOVOT -Israel 'DSM - GELEEN - HoI/and

Using polymeric fibre strength data from the literature and from our own laboratory tests, we show in the first part of the present paper that simple modifications of 1'0isson/Weibull concepts, resulting in a new failure probability function, can he used as a modeling scheme for the study of diameter eflects 011

strength in a way at least as satisfactory as previously used LEFM-based schemes. In the second part of the paper a new experimental approach for the study of composite failure is presented. Specially prepared composite monolaycr models were tested in simple tension under an optical microscope equipped with crossed polarizers <lnd with a video camera. The potential usefulness of this approach for the characterisation of hasic failure modes in fihre-reinforced composites, and as a probe of existing strength theories in such materials, is demonstrated.

INTRODucnON

An accepted approach to the problem of fracture nucleation in unidirectional composites is to view it as a complex statistical/stochastic process involving scattered failure of fibers at flaw sites, and local overloading and failure of neighhouring fibers by way of stress transfer through the matrix. It is also commonly predicted, via several variants of a model for crack growth, that final (catastrophic) failure follows via the rapid development of a cluster of neighbouring broken fibers, possessing a critical dimension. The Poisson/Weibull probahilistic approach is the most popular scheme in use for describing the ultimate mechanical hehaviour of solids having a linear stress-strain behaviour up to breakdown. It has also been used to describe other types of extreme-value prohlems /I, 2;' The effects of specimen dimcnsions, termed size effects, are of great importance in the strength characteristics and are inherent to the Poisson/Weihull/weakest link modcl, hoth for a fibcr and a fihrous composite. This is discussed in details elsewhere /I, V Regarding a fiber the length and diameter are known to he key factors with respect to the mechanical stress

72

necessary to promqte failure, and different methods are usually utilized to quantify both types of such size effects: linear elastic fracture mechanics (LEFM) and related schemes provide the theoretical basis for the effect of diameter variability upon strength whereas statistical theories, generally based upon the Weibull probability distribution combined with the weakest-link theorem, describe length effects. In fact size effects of both types (length and diameter effects) can be conveniently quantified by a unified statistical approach, as shown here. Regarding a composite size effects are also included in theoretical expressions for the strength under a linearly increasing load history and for the lifetime under a constant load history. Experiments on this, however, are scarce. Here we describe a new experimental approach to generate preliminary results on the elastic and fracture properties of several types of composite monolayer models, including hybrids. Future tests will deal with size effects in such models as well.

I - EFFECT Of FIBER DIMENSION ON STRENGTH

1.1. The tEfM model

Linear elastic fracture mechanics (tEFM) and related schemes usually provide the theoretical basis for the effect of fiber diameter variability I1pon strength through a relationship of the form /4, 5/:

(I)

where cr is the strength and D is the fiber diameter. This is a Griffith-like expression but in which the fiber diameter rather than the critical crack length is used.

1.2. The statistical/stochastic model

Statistical theories, generally based upon the Weibull probability distribution combined with the weakest link concept, describe length effect via the expression /6/:

(j' '" r lib (2)

where 0 is the mean strength at a given (nondimensional) length I, and b is the shape parameter in the two-parameter Weibull probability distribution for strength.

1.3. The modified Poisson/Weibull scheme for diameter effects

Size effects of both types (length and diameter) may in fact be conveniently quantified by a unified statistical approach. First, it is clear that Wei bull plots can be performed correctly only if the strength data are generated from a population of geometrically identical specimens (precisely because size effects are known to occur). One way to take fiber to fiber diameter variability into account is to use the following failure distribution function /1/:

F(cr) = I -exp{ -kd~ crb} (3)

73

where F(cr) is the probability of failure of a fiber under a stress cr or less, d is the (nondimensional fiber diameter, k is a positive constant, and 1) is a real constant. With the failure distribution function given by equation 3, termed here the modified Poisson/Weibull distribution, the slope of a In (0') vs. In (d) plot is equal to -Sib, rather than to -lib as in a conventional Weibull model. As seen, the shape parameter is not anymore the only parameter determining the size effect as in previous statistical model: rather, -Il/ b is now the determining factor. At the condition that 1) takes on the appropriate numerical value, the values of b obtained from a modified Weibull plot -through linearization of equation 3- and from the In(strength) vs. In(diameter) plot are reconciliated.

1.4. Experimental evidence

Recent experimental work II, 2, 7, 81 amply justifies the approach proposed in section 1.3. As an example, four sets of ultrahigh strength polyethylene (UHSPE) fibers were used 121 and shown to fit better equation 3 than a conventional Weibull model (see figure I). Moreover, the modified Poisson/Weibull approach was shown to fit a In(strength) vs. In(diameter) plot in a way at least as satisfactory as a LEFM-like model. Finally, the modified statistical model reconciliated the values of the shape parameter b as calculated from a Weibull plot and a In(strength) vs. In(diameter) plot. Various other examples are availahle in the above cited references.

2 - DAMAGE ANALYSIS IN COMPOSITE MONOLAYERS

2.1. Materials and methods

Composite monolayer models, which consist of carefully positioned single fibers into a thin epoxy matrix layer, were fabricated using a specially developed technique. Details of the manufacturing method are given elsewhere /3, 9/. The monolayers, of various types, were tested to failure in simple tension using a custom-made minitensile testing device fitted to the stage of a Stereozoom microscope. The fracture nucleation and growth process was followed by video microphotography. The experimental setup is sketched in figure 2. The fiber used were E-glass (Vetrotex), Kevlar 29, Kevlar 49, and Kevlar 149 (du Pont). The matrix system was CY223/HY956 epoxy (Ciha Geigy). As an example, a composite monolayer is presented in Figure 3.

2.2. Fracture results: preliminary conclusions

Complete details of the results regarding our preliminary tests are given elsewhere /9(. As an example in Figure 4 the modulus of Kevlar 149/epoxy monolayers made of I single fiber, and of 8 single fibers, is seen to follow quite well the Halpin-Tsai (or rule-of-mixtures, RoM) equation, even at such low volume fraction'. Also, the strength is shown to increase linearly according to the RoM. These results, although preliminary, are encouraging and confirm that the quality of the manufactured microcomposites is high. The study of microcomposites via videomicrophotographic techniques is particularly relevant to the dynamics of mechanical fracture in such materials. Indeed, with respect to failure dynamics, the first findings of our research are as follows: (i) As predicted by several authors (see references in /9!), a critical number of neighbouring fihres must break hefore fast failure occurs. Based on a very limited number of preliminary tests, this number was found to be equal to 4 in E-

74

glass/epoxy microcomposites. In Kevlar 149/epoxy microcomposites, the crack growth process is extremely fast, but some fibre integrity is conserved (fibre splitting rather than snapping occurs). (ii) In principle information on the type of load sharing rule in effect in a given fibre/matrix system system can be obtained. A fast­propagating damage array is observed to nucleate and grow in Kevlar 149/epoxy systems, at several places simultaneously. This may indicate that equal-load sharing is in effect in this system. (iii) In principle, the effect of changes in the chemical or physical nature of interfaces on the failure nucleation, growth, and criticality can be studied. This is currently pursued. (iv) The effect of fibre bunching (that is, the fibre­to-fibre interaction) and/or of volume fraction, as well as of fibre misalignement, on the mechanical and fracture properties can easily be assessed. This is also currently being studied.

ACKNOWLEDGEMENTS

IT. D. Wagner is the recipient of the .I. and A. Laniado Career Development Chair. This work was supported by the National Council for Research and Development, Isracl, and the K. F. A. .luelich, Germany.

REFERENCES

I. II. D. Wagner,.I. Polymer Science - Polymer Physics, in press (1988).

2. H. D. Wagner, L. W. Steenbakkers, Phil. Mag. Letters, in press (to appear in 1989).

3. II. D. Wagner, in Application of Fracture Mechanics to Composite Materials (R. B. Pipes, K. Friedrich, Eds.), Elsevier Sc. Publ. B. Y., in press (to appear in 1989).

4. C. Gaiiotis, R . .I. Young, Polymer 24, (1983) 1023-1030.

5. J. Smook, W. Hamersma, A . .I. Pennings, .T. Materials Science 19 (1984) 1359-1375.

6. W. Weibull,.T. Applied Mechanics 73 (1951) 293-299.

7. L. W. Steenbakkers, H. D. Wagner, J. Materials Science Letters, in press (to appear in 1989).

8. A. S. Taylor, II. D. Wagner, submitted paper (1988).

9. H. D. Wagner, L. W. Steenbakkers, submitted paper (1988).

75

(0)

1 1 SPECTRA 1000 SPECTRA 1000

WEIBUU. MODIFIED WEIBULL

N = 32 N = 32 a 0

a = 2.79 GPa a = 2.55 GPa

b = 2.96 b = 6.3

oS = 12.2 ~ -1 -1 r: I ...

3' I -2 '5 0

0

-3 0

-4

-1 -0.5 0 0.5 1 1.5 2 -1 -0.5 0 0.5 1 1.5 2

LN(Strength. GPa) LN(Strength. GPa)

Fig. I - Weibull plots of LJIISPE Spectra 1000 strength data using (a) the conventional Weibull model and (h) the modified model hased on equation 3.

76

,

, .. ·· ... · .......... ······· ... ··············IM~"I

Fig. 2 - The experimental setup, including a PS/2 microcomputer and a video graphic printer currently added.

--y . t ,;. ,

It . 4'" -

Fig. 3 - SEM micrograph or a Kevlar 49/epoxy microcomposite monolayer after ~ rracture test (har is 100 ~Im).

-;;;

o = Epoxy film (95,. c.i.)

10 X = I fibre monolayer

o = 8 fibre monolayer

0.125

"BRE CONTENT BY VOLUIIE, Vf

(b)

o = Bpoxy mm (95,. c.l.) I I x = I fibre monolayer KEVUR J(9!cY223-HY958

o = 8 fibre monolayer

~ 0.100

II

~ 0.075 Z .. It: t; 0.050

0.025

0.000 0.005 0.025

FIBRE CONTENT BY VOLUIIE, V r

77

fig. 4 - (a) Young's modulus and (h) tensile strength against fiher content 111

Kevlar 49/epoxy microcomposite monolayers.

ABSTRACT

THE STRENGTH OF TUNGSTEN·CORED SILICON·CARBIDE FIBRES

AND THE INFLUENCE OF A POLYMER MATRIX

M.G. BADER', D.A. CLARKE"

'University of Surrey - GUILDFORD - England "Rolls Royce PLC - DERBY - England

Silicon carbide fibres of 100 pm diameter were tested in tension at various gauge lengths in air, embedded in two polymer matrices and with thin coatings of a polymer. It has been shown that the tensile strength distribution is strongly influenced by these variables. It appears that coating or embedding inhibits failure from surface flaws at strains which would otherwise have led to failure.

1-INTRODUCTION

The basic objective of this work was to study the effect of a matrix on the tensile strength of a fibre. The tensile strength of the selected fibre was measured by testing single fibres in air and then embedded in two resin matrices. This showed a significant strength enhancement for the embedded fibres, so further tests were conducted on fibres coated with thin layers of the matrix resin.

2-EXPERIMENTAL

The fibre used was taken from a single length of 100 pm diameter, tungsten-cored silicon-carbide which had been manufactured by the chemical vapour deposition process (CVD). Conventional tensile tests were carried out on single lengths of fibre of from 10-500 mm gauge length; 30-90 tests were conducted at each of 4 lengths and the data analysed by the maximum likelihood method to determine the Weibull characteristic strengths and exponents.

Lengths of fibre were then embedded into resin coupons and these were tested in tension. Further fibres were coated with thin layers of resin and then tested.

80

It should be noted that in the case of the fibres tested in air and the coated fibres, the maximum force at failure was measured and the stress was calculated from the measured cross sectional area. However, from the embedded fibres the strain at failure was measured and the stress calculated fr~m a knowledge of the Young's modulus which had also been measured. In the tables of results both stress and strain at failure are given and the measured parameter indicated by * 3-RESULTS

3.1 Single fibres tested in air

The data from these tests are given in Table 1. Tests were conducted at 4 gauge lengths and these show the expected trend towards higher mean strengths at the shorter lengths. The characteristic failure strains as plotted in Fig 1 show a superficial conformity with the Weibull 2-parareter model (ie linear) , however, a more detailed analysis has revealed considerable discrepancy.

This is due to the fibres failing in two distinctly different modes The low strain failures appear to initiate from surface flaws~ whilst at high strain, failure emanates from the centre, probably following yielding or failure of the tungsten core. Tests on the tungsten core-wire showed it to yield at a similar strain. It is likely, however, that within the fibre the tungsten would have become embrittled during the CVD process and fracture rather than yield would occur at this strain.

3.2 Single embedded fibres

For these tests single fibres were cast into 4 mm thick sheets of resin. Coupons 20 mm wide and 150 mm long were cut from these sheets with a fibre located along the axis. These were extended in tension and failure of the fibre observed by viewing through crossed polarisers. The fibre break resulted in a local photoelastic disturbance in the resin which rendered it easily visible. Two resins were used, a polymethylmethacrylate (PMMA) polymerised in situ from the monomer and a plasticised anhydride cured epoxy. The strain on the coupon was monitored and strains at fibre failure were recorded. The results of these tests are given in Table 2 and are also shown in Fig 1. It will be noted that the embedded fibres fail at a much higher strain than fibres tested in the air at similar gauge lengths. For instance, at 50 mm gauge length, the PMMA embedded fibre fails at 1.23% strain, the epoxy embedded fibre at 1.08% whilst the unsupported fibre failed at 0.71%. These are all Weibull characterisitc strains estimated by maximum likelihood from 13 - 90 tests. Part of the increase may be attributed to compressive strains which are induced in the ffbre by shrinkage of the resin during cure and on cooling to room temperature after cure. This would be greater for the PMMA as it had been polymerised directly from the monomer. Many broken fibres were extracted from these coupons and, without exception, showed core-initiated fractures whereas fibres tested in air showed nearly all surface-initiated failures. It, thus, appears that when embedded the surface flaws are somehow inhibited.

81

3.3 Coated Fibres

In order to test this hypothesis a further batch of fibres were prepared with thin coatings of epoxy resin deposited on their surface. Coatings with thickness of 5,30 and 40 }lIII were used. These fibres were then tested in air and the results are also given in Table 2. These show a progressive increase in characteristic failure strain from 0.92' for the 5}l111 coating up to 1.05' at 40}JIII which corresponds closely to the value recorded for the fully embedded fibre. It was not possible to collect the failed fibres with the 5 pm coating for examination but some 50' of those from the 30 and 40 pm coated fibres were examined in the SEM and found to be core initiated.

3. 4 Failure mode

In addition to the observation of core and surface failure in the fibres it was observed that in the fully embedded test pieces there were indications of ! complex stress redistribution either side of the point of failure 'Fig 2. Initially this was thought to be due to debonding but several coupons were sectioned so as to expose the fibre, Fig 3, and it was seen that multiple failure of the fibre had occurred with fractures typically 200 pm (2 diameters) apart. When the resin was dissolved away these fragments could be seen, Fig 4. The explanation of this phenomenon is that the fibres adhere only weakly to the matrix and when the initial fracture occurs the energy release is sufficient to set up a shock wave which leads to a number of secondary tensile fractures. All the fractures appear to have been core initiated. This phenomenon is considered to be more likely to occur in these larger diameter fibres where the surface area to volume ratio is much smaller than in the more usual glass and carbon fibre systems. The fact that the fibres are this size also renders the effect much easier to observe.

4-Concluding remarks

The principal conclusion of this work is that embedding a fibre in resin or even merely applying a thin coating may have a significant effect on the strength distribution and mode of failure. In this case it appears that surface flaws which control the lower part of the strength distribution have been rendered ineffective and the strength thereby enhanced. This will not necessarily happen to the same extent in other systems but the possibility must be considered when using data gathered from single fibre tests to predict composite behaviour.

5-REFERENCES

1. D A Clarke, PhD TheSis, University of Surrey 1988 "The strength of model composites incorporating silicon carbide fibres".

2. D A Clarke and M G Bader in Proc ICCH-VI and ECCM-2 Ed F A Matthews et aI, (1987) 5.382-5.392, Elsevier, London.

3. D A Clarke and M G Bader Mater Sci Letters, (1986), 903-904.

82

RESULTS

TABLE 1

Tensile strength of 100 ua silicon carbide fibres

Gauge No. Neibull Neibull Neibull length tested Characteristic Characteristic exponent

(mm) failure strain failure stress n Eo 0'0 W

(\) (GPa)

10 71 0.75 2.60* 2.9

50 90 0.71 2.47* 3.6

100 30 0.60 2.08* 4.0

500 50 0.55 1.91* 3.2

TABLE 2

Tensile strength of silicon carbide fibres at 50 .. gauge lengths

n EO 0'0 W Fibres in air 90 0.71 2.47* 3.6

Fibres in PMAA 36 1.23* 4.27 7.7

Fibres in Epoxy 13 1.08* 3.75 10.2

Coated fibres

5pm epoxy 39 0.92 3.19* 4.9

30pmepoxy 46 0.98 3.40* 8.6

40pm epoxy 30 1.05 3.65* 6.4

see Table 1 for explanation of symbols

* measured parameter

83

fibre length 20mm 100mm 500mm

10~m I 50mm I 200;nm

0.4 ~ ..--.. 0.2 ~ '--"

c 0.0 0

-0.2 L.. -III Q) -0.4 L.. :J -0.6 0

'>--O.B c

0 -1.0 Q)

* E Fibres embedded in epoxy matrix "--' -1.2 c 0 Fibres embedded in PMMA matrix -1 -1.4 • Unsupported single fibres

- 1.6 -6 -5 -4 -3 -2 - 1 0.0

Ln [fibre length (m)]

Fig. 1 Plot of In characteristic failure stress vs In fibre length for silicon carbide fibres tested in air and embedded in two resin matrices.

Fig. 2 Transmission optical micrograph under crossed polarizers of the zone around a fibre failure for a silicon carbide fibre in an epoxy resin matrix. The bright zone at the extreme left-hand side marks the point of stress transfer back into the fibre.

84

Fig. 3 An optical micrograph of a polished section through a similar fibre fracture zone to that depicted in Fig. 2. The tungsten core and multiple fracture are clearly shown.

Fig. 4 Fibre fragments rema~n~ng after the matrix resin was dissolved away from a fractured embedded fibre similar to those shown in Figs. 2 and 3.

MATRICES CERAMIQUES CERAMIC MATRIX

Chairmen: Mr B. BRoaUERE Societe Europeenne de Propulsion Pr B. HARRIS University of Bath

FIBRE REINFORCED ALUMINA CERAMIC COMPOSITES BY SOL-GEL PROCESSING

ABSTRACT

M. CHEN, P.F. JAMES, F.R. JONES, J.F. BAILEY

University of Sheffield School of Materials, Northumberland road

SHEFFIELD S10 2TZ - England

The sol-gel route to unidirectional fibre reinforced ceramic composites has been demonstrated. The technique has been modified to increase the solid yield of the sol using a particulate filler. After sintering the composites were found to be microcracked as a result of the constrained sintering and thermal shrinkage. The interfacial shear strength has been estimated from the average spacing of the microcracks, and found to be comparatively low for the carbon fibre reinforced composites.

1. INTRODUCTION

The sol-gel processing route to ceramics, which has recently been developed enables a relatively low sintering temperature to be employed. This technique has bene examined for the production of fibre reinforced ceramic composites (FRC) because of the potential microstructural uniformity of the matrix and the ease of fibre dispersion. The sol-gel route to alumina ceramics developed by Yoldas (1) in 1975 is widely adopted, but for FRC multiple impregnation of fibres is usually required (2) to counteract the low solid yield from the sol. Therefore, the present work has focussed on methods of increasing the solid yield of an alumina sol whilst maintaining the processing benefits of the sol-gel technique and on the characterisation of the fabricated composites.

2. EXPERIMENTAL

Boehmite gel powder, containing 0.2% Ti02 was used as the precursor for the sol and 0.5 ~m (Alcoa) a-alumina particles were used

88

as the filler. The alumina fUler was deflocculated and dispersed 1n aqueous nitric acid solution at a pH of 3.5 and a volume fraction of O.~ and the boehmite sol was peptized with HNO", at a pH of 3.5 and a volume fraction which will give a solid yield of ex-A1 2 0", of 10~ by volume (v%). Sols with differing filler content were prepared by mixing the two sols. The filler content is reported as a volume fraction of the final alumina ceramic assuming full sintering to theoretical density. The sol was gelled by adding dilute aqueous aluminium nitrate.

Unsized surface treated high strength carbon fibres (Grafil XAS) and Nicalon SiC fibres (ceramic grade from Nippon Carbon) were used as the reinforcements. The carbon fibres were first coated with a thin layer of boehmite gel and fired to 500'C for 0.5 hr. The SiC fibres were fired to 600'C for 5 min to remove the sizing. Fibres were aligned in a rectangular mould and infiltrated with a sol containing 80% filler and 20% boehmite gel. The matrix for the composite was gelled in 20 min after removal of trapped air.

The gels and the unsintered composites were dried at 50'C for 20 hr and heated slowly in N2 to 500'C and then at 5 K min- 1 to 1000'C and at 10 K min- 1 to the sintering temperature where the samples were held for 2 hr and finally allowed to cool overnight. Composite samples were also uniaxially hot pressed at 1200'C for 1 hr, using a graphite die. To identify the cracking in the matrix CF composites were sintered in N2 followed by exposure to air at 100'C below the sintering temperature for 6 hr to burn out the fibres. The porous matrix was then cooled down, and impregnated with mounting resin and polished. The density and the micro-hardness of the I118trix and the flexural strengths under three point bend were determined.

3. RESULTS AND DISCUSSION

3. 1. Matrices

The modified sol-gel system has been described in detail elsewhere (3). The total solid yield of alumina of the sol was increased as the filler content was increasd without causing immediate gelation. Thus the solid yield of ex-alumina could be increased from 10 v% for the unfilled sol to 32 v% with 90% filler and a solid yield of 25 v% for the sol containing 80% filler. It was shown that after addition of the gelling aid, Al (NO",)", there was a period of about 10 min during which the viscosity remained sufficiently low for infiltration prior to gelation.

After drying at up to 500'C all the molecular and structural water in the gel was removed, giving a green density which was 35% of theoretical for ex-alumina. With 80% filler it could be increased to 58%. After sintering at 1200'C for 2 hr the unfilled gel had a density 60% of theoretical whereas the 90% density was reached .with a gel containing 10% filler (Fig. 1>. At higher filler contents the enhancement was smaller but with 80% filler 69% density could be obtained. Further sintering to 1450'C yielded 95% density, with a uniform microstructure and a grain size of <5 ~n

89

The acidic peptisation process leads to positively charged particles in the boehmite sol and the a-A1 2 03 dispersion and a stable mixed sol of low viscosity. Since the alumina filler is larger in particle size and denser than the boehmite colloids it was possible to disperse it to a higher solids content and increase the solid yield. As the boehmite colloids floculated to a gel, a rigid continuous gel matrix was formed holding the dispersed filler together. After drying the green density (58\) was lower than that of the slip cast sample (65%) because of incomplete dispersion and low packing efficiency of the filler particles in the presence of the gel colloids. Becher et' 151. (4) showed that AIOOH gel is transformed into a porous a-Al"O" having a similar sintering behaviour to the cold pressed powder compact. However, in the, presence of the fine a-Al.,O:. particles the phase transformation is nucleated heterogenerously to give a refined microstructure of high sinterability (5). However for the highly filled gel, although a sillilar process will occur, sintering will be inhibited when the filler particles are in contact.

3.2. CompOSites

For the carbon fibre composites a uniform fibre dispersion was readily achieved by this fabrication route. In the case of SiC fibre composites the fibre tows are less easily dispersed,

The green density of the SiC reinforced composites, with a fibre volume fraction of 0.31, is 67.4% of the theoretical fully dens1fied composite (calculated from the rule of mixtures). This can be increased to 72.5% by sintering at 1200'C and further to 74.3% with hot pressing. In the absence of fibres the green density of 58% was increased to 69% after firing. From the measured density and Vf of the composites the density of the matrix was found to have been increased from 55% to 59%. Since this is lower than that of matrix in the absence of fibres it implies that the shrinkage of the matrix during sintering is inhibited by the fibres. This was confirmed by microhardness measurements where in the composites, values of 3.28 -7.58 GPa were found. The matrix alone had values of 4.17 - 9.0 GPa. Although the microhardness is lower than that of the corresponding monolithic alumina it is increased with sintering temperature. When sintered to 1200'C the microhardness of matrices ranged 3-4 GPa. Similar results were reported by Coyle et 151. (6).

In order to confirm this effect careful microscopy has been carried out. As shown in Fig. 2, after the crack free SiC-composite was sintered to 1200'C evenly spaced transverse microcracks were observed with average spacing of 42 ~m. This indicates the formation of thermal cracks during sintering and cooling which arise from the induction of residual tensile stresses. These stresses could arise from shrinkage during sintering and from a mismatch in thermal expansion coefficients. Considering the residual stresses in the matrix in the longitudinal direction parallel to the fibres it is clear that cracking will occur when the matrix strength is less than the sum of the thermal and sintering components. Attempts have been made to identify the mechanism of Ilicrocracking. To determine if matrix cracking occurs during sintering of the carbon fibre composite

90

the fibres were burnt out at- 100·C below the sintering temperature of 1200·C. As shown in Fig. 3 cracks were already present at an average spacing of 100 ~n The cooling interval of 100·C is considered to be insufficient to cause cracks at this denSity. This is confirmed since at the higher sintedng temperature of I~OO·C the microcracking was found to be more severe. Furthermore on cooling in the absence of fibres, therllal cracking should be absent. It- is concluded, therefore, that these cracks must result from the sint ering process. Since the crack densities in this case are apparently similar, it suggests that saturation cracking had occurred under the influence of the &intering stresses in this case.

Cooper and Sillwood (7) have demonstrated that the sliding frictional stress ('t) acting at the fibre-matrix interface could be evaluated from the evenly spaced microcracks, according to equation (1)

't = V", (]",ur IV~ 2x' (1)

where (] .. u is the ultimate strength of matrix, r is the diameter of fibres x' is the minimum crack spacing. It has been demonstrated statistically that x' = xiI. 3~ where x is the average crack spacing. If we assume that saturation cracking had occurred during fabrication then this model can be used to estimate 'to

In this experiment, the SiC fibres and carbon have average diameters of 15 ~m and 8 ~m respectively. The average V, for the composites were 0.33 and 0.42 respectively. The value of (]mu was taken to be that of the equivalent monolithic alumina, which was found to be 109.5 MFa under three point benidng. The tensile fracture stress of the matrix is expected to be lower than that of the monolithic material. However the flexural strength is considered to be a good aproximation for (]mu' The average crack spacing was observed to be 42 Jim and 100 ~m, from which 't is estimated to be 55.9 and 8. 1 MFa for SiC and C composites respectively. The former 1s in good agreement with the value of 47 MFa reported by Coyle et al. (6) for slip-cast composites.

Mechanical performance of the as-prepared composites under three point bending are illustrated in Fig.4 by the load-deflection curves. The flexural strength of the carbon fibre reinforced-deflection alumina was 345 MFa and the flexural modulus Eo: = 95.9 GPa. For the SiC composite values were 324 MPa and E", = 66 GPa respectively. As shown, the composites fail in a non-brittle manner by multiple matrix fracture. They were still able to carry a significant load, even though they were microcracked. During testing the carbon fibre composite failed by shear, confirming the low interfacial shear strength indicated by the microcrack density. The theoretical estimates from the law of mixtures can be calculated assuming that the matrix is microcracked and non-load bearing and to a first approximation, E" = E~V~. Thus for the SiC composites where (]fu = 1000 MPa, E~ = 200 GPa (6,8) and (]mu = 109.5 MFa, it is estimated at a V~ = 0.33 that (]", = ~06 MPa and Eo: = 66 GPa. These agree well with the experimental results and confirm the micromechanical model. For the C-fibre composite taking (]~u = 4000 MPa, E, = 225 GPa and Vf

0.~2, it can be calculated that (]c = 1740 MPa, Eo: = 96 GPa. The actual (]eu is far lower.

91

4. CONCLUSION

It has been demonstrated that unidirectional fibre composites can be prepared via the modified sol-gel process with one single infiltration. After sintering the composites preform, a partially densified and cracked composite was obtained. These cracks were induced both by therJlal and sintering shrinkage. The crack densities illustrated the relatively poor interfacial bond achieved with carbon fibres which was also reflected in the stress-strain response of the composite.

10'".----- ---------- , ...... ~ .... 9

80

~ ~ 70 .~

o 60

~ :c: ~50 II:

~'

, ,

.~

/ /

~;f- ' ,

Healing ru(. 2II"C1"~n I-Iold lor 2 hr

~~

. ..... .. _ ... ..

•

• / , ,

-' -'

FIII.t co,,1 0 nlt'/,)

• 0 • 10 • 80

Fig.l gels

Relative density with various

contents at different FIring remp.ralur. I"C' temperatures.

Fig.2 Polished surface of SiC/A1203 composite showing regularly spaced microcracks after firing at 1200°C for 2hrs in N2 Vf=O.33.

of fired filler firing

92

Fi¢.:l Polished surflH't's of C compositt's showing regularly spaced microcracks aftt'r firin¢ at 1200 C in N2; (11) natural cooling in N2'

(b) after .,XPOSUl't' ttl air at 1100'C prior to cooling. (bars donate 100 pili)

\J) \J)

~ \J)

I I ceFlECTI~ mm

300

200

100

Fi~.~ Typical fracture behaviour of ,'einforced alumina composites using

REFERENCES

carbon three

fibre point

i I ! ~h 'Cl---::r

h b l 1,19 5] 31mm ;i :1.7 58 l,Orrm

S iC/A1203

and SiC bending.

fibre

1 - Voldas, B.E., Am. Ceram. Soc. Bull., 54(3) (1975) 268-88. 2 - Fitzer, E. and Gadow, R., Conference on Tailoring Mult1phase

and Composite Ceramics, Penn. State Univ., July 1985. 3 - Chen, M., Bailey, J.E., James, P.F. and Jones, F.R., lnst.

Phys. Conf. Ser. No. 89: Session 5 (Warwick, 22-25 September 1987) p.183.

4 - Becher, P.F., Sommers, J.H., Bender, B.A. and MacFarlane, B.A., Mat. Sci. Res. Vol. 11 (H. Palmour et al. eds.) p.79 (New York: Plenum) .

5 - Kumagai ~ and Messing, G.L., J. Am. Ceram. Soc. 68(1985) 500. 6 - Coyle, T.W., Guyot, ~ H. and Jamet, J. F., ceram. Eng. & Sci.

Proe. 7(7/8) (1986) 947-57. 7 - Cooper, G.A. and S11lwood, J.M., J. Mat. Sci. 7 (1972) 325-33. 8 - Clark, T.J., Arons, R.M. and Stamtoff, J.B., Ceram. Eng. & Sci.

Proc. (7-8) (1985) 576-88.

ABSTRACT

SILICON CARBO·NITRIDE CERAMIC MATRIX COMPOSITES BY POLYMER PYROLYSIS

R. LUNDBERG, P. GOURSAr

Swedish Institute For Silicate Research Box 5403 - 40229 GOTEBORG - Sweden

·Laboratoire de Ceramiques Nouvelles, LA CNRS 320 UER des Sciences Albert Thomas - LIMOGES - France

Polymer pyrolysis as a route to fibre reinforced Si-N-C materials was identified and demonstrated. The synthesis and pyrolysis of polysilazane precursors were optimized with respect to ceramic yield and minimum bloating during pyrolysis. Solution infiltration of stacked Nicalon SiC fibre weaves was performed. Repeated impregnations/pyrolyses yielded fairly strong composites exhibiting non-brittle fracture behaviour, especially when a polymer free from low molecular weight oligomers was used and the fibres given a carbon coating.

INTRODUCTION

Silicon nitride is one of the most promlsmg high performance ceramics for engineering applications such as heat engine and gas turbine components, high temperature heat exchangers and wear parts. A non-catastrophic "graceful" failure with controlled crack extension would significantly increase the possibility of using Si3N 4 for many applications. An obvious way to achieve this is through the composite approach. Several other ceramic matrices have been successfully reinforced with continuous ceramic fibres. Glass ceramics /1/ and chemical vapour deposited (CVD), or rather infiltrated (CVI), SiC /2/ matrix composites are today being introduced in various applications where their non-brittle fracture behaviour coupled with good high temperature properties make them the material of choice. Despite an early interest in Si3N 4 matrix composites /3//4/ their development has not been as successful as for the aoove mentioned matrices. The high sintering temperature required for Si N as compared to glass ceramics and the CVD temperature for SiC has made progr~ss 41ess straightforward. Existing ceramic fibres, such as the polymer derived Si-C-O Nicalon fibres are limited to processing temperatures in the range of 1000 - 1200 °C. An interesting way of producing a silicon nitride matrix at lower temperatures is polymer pyrolysis. A fibre preform is infiltrated with molten or dissolved polymer and subsequently pyrolysed, typically around 1000oC, to yield a ceramic residue. This method is today used extensively and successfully to produce carbon/carbon composites /5/ and has been demonstrated for SiC/SiC composites using polycarbosilanes /6/!7 /.

94

The synthesis and pyrolysis of polysilazanes as precursors for Si3N 4 or Si-N-C ceramics has been studied recently /8//9/. Polysilazanes have been used as polymer precursors for a new Si-N-C ceramic fibre /9/. In the present investigation a cyclic silazane is polymerized and the polymers obtained are used as matrix precursor for fibre reinforced silicon carbo-nitride matrix composites.

I - EXPERIMENTAL PROCEDURE

1.1. Processing

Polycondensation of octamethylcyclotetrasilazane (OMCTS) with 2.5 wt-% KOCH3 as catalyst was carried out under reflux conditions in nitrogen at 3600C /10/ for several hours. Some polymers were post-treated in flowing N at the polymerisation temperature to remove low molecular weight oligomers. Fibre ~reforms in the form of stacked weaves of SiC-fibres (Nicalon NLM-I02, Nippon Carbon CO.,Ltd., Japan) were infiltrated with polymer dissolved in tetrahydrofuran (THF) and dried under light uniaxial pressure. Pyrolysis of the infiltrated preforms was performed in nitrogen atmosphere, typically following a temperature-time cycle as shown in Figure I, reaching the top temperature (between 900 and 1200 0q in about 30 hours. After the first pyrolysis several impregnations could be made without pyrolysis in between. The solution dried Quickly leaving solid polymer in the pores that did not dissolve completely during the next infiltration. In this way the porosity was entirely filled with polymer before the second pyrolysis. Some samples were produced with a carbon coating on the fibres. This was obtained by dipping the weaves in a THF solution of a pheno~c resin (Peracit 5046, Perstorp Chemitec, Sweden) followed by pyrolysis in N2 to 900 C.

1.2. Testing

Thermogravimetric analysis (TGA) was performed in nitrogen with a linear temperature increase (750 °C/h). Fracture behaviour of the pyrolysed composites was studied in three-point bending on as-pyrolysed samples, thickness 2 mm, span 22 mm at a cross head speed of 0.5 mm/min. Microstructural features were examined with optical microscopy and scanning electron microscopy (SEM). Outer surfaces, fracture surfaces and interlaminar surfaces (as obtained by deliberate delamination) were studied on pyrolysed composites using a Jeol JSM-35 SEM.

II - RESULTS AND DISCUSSION

2.1. Polymer pyrolysis

The polymerisation reaction, the influence of temperature and catalyst content on structure and polymer yield as studied with IR spectroscopy, gel permeation chromatography, mass spectrometry and small angle x-ray scattering (SAXS) has been described previously /10//11//12/. The polysilazanes were found to consist of ordered microdomains separated from each other by low density areas. These particles made of stacked polycyclic macromolecules or folded sheets are linked by chains or oligomers. They are white solids highly soluble in organic solvents (such as THF).

During the pyrolysis the polymers are gradually converted to amorphous Si-N-C ceramics, with evolution of CH4, H and some NH . The thermal decomposition proceeds in three stages, as shown in tigure 2. The firsl stage (250-450 Oq is related to the decomposition of the low molecular weight oligomers. This first stage, i.e. the content of oligomers, is important for the microstructure of the resulting ceramic. Polymers with a high oligomer content exhibit problems with bloating. A large gas

95

volume is evolved while the polymer is still viscous resulting in bubble formation. The polymer solidifies before the bubbles disappear and a porous ceramic foam structure is obtained. Polymers treated in flowing nitrogen at the polymerisation temperature to remove the oligomers were found to be significantly less sensitive to bloating. After pyrolysis (as shown in Figure I) of such a post- treated polymer, dense bodies were obtained. During the second stage, between 450 and 750 °C, the decomposition behaviour is very similar in spite of differences in oligomer content. Finally, between . 0 750 and 1200 C the pyrolysis products are stable.

2.2. Composite microstructure

Microstructural investigation of the composites after pyrolysis showed that the polymer solution had infiltrated even thick preforms. However, the difference in pore size between interlaminar pores, pores between the strands in the weaves and the very fine porosity within the strandsresulted in non-uniform matrix deposition. The smaller pores were filled first, sealing off the large pores between the weaves from the surface before full density was reached. This created defects (pores) of several hundred I'm size. However, the composites retained their shape after pyrolysis with no evidence of bloating or delamination, even though the poor wetting and film forming properties of the polymers resulted in an inhomogeneous matrix with drying cracks and pores as large defects. These problems might be reduced if the first infiltration was done by melt infiltration to more effectively fill the large pores and eliminate the need fo.r drying. The micrographs in Figure 3 show large cracks and pores on the outer· surface (3a) and the good infiltration of the matrix within the strands seen on a fracture surface (3b).

2.3. Fracture behaviour

So far only limited mechanical testing has been performed indicating a fairly strong fibre/matrix bond although no chemical interaction could be seen between fibres and matrix. The material produced with the polymers having a high oligomer content was very weak, with large processing defects but still fractured in a non-brittle manner (see Figure 4). When using the oligomer-free polymers and a carbon coating on the fibres a more homogeneous material with non-brittle fracture behaviour and a maximum stress level of around SO MPa was obtained (see Figure 4). This material stilI contains almost mm size defects and large voids, especially between the weaves, and it could be expected that better processing, such as by melt infiltration, would significantly increase the strength level obtained. The non-brittle fracture behaviour can be seen on the load/displacement curves in Figure 4, as well as on the fracture surface in Figure 3b. The rough fracture surface with significant fibre pull-out is characteristic of a high toughness composite.

CONCLUSIONS

Polysilazane solution infiltration of a SiC (Nicalon) fibre preform with subsequent pyrolysis was shown to be a promising route to fibre reinforced silicon carbo-nitride. Pyr0l1.sis could be carried out, without bloating of the polymer, at temperatures around 1000 C i.e. were the fibres are not thermally degraded. Inhomogeneous polymer deposition and poor film forming properties of the polymers led to composites with large defects in the form of drying cracks and pores. With a carbon coating on the fibres and a polysilazane optimized with respect to bloating, composites failing in a non-brittle way with a maximum stress level of around 50 MPa were obtained. Further development in polymer precursor synthesis and infiltration processing (melt infiltration) is needed to obtain high strength, dense, flaw-free composites. In view of

96

the success of carbon/carbon materials /5/ polymer infiltration and pyrolysis has a large potential as a method of producing ceramic/ceramic composites.

ACKNOWLEDGEMENTS

The Swedish National Board for Technical Development (STU) is thanked for the scholarship financing the one year stay of R. Lundberg at Lab. Ceramiques Nouvelles, Limoges, France, as is Prof. M. Billy for making this stay such a fruitful experience. The authors also wish to thank Mr. L. Eklund and Miss C. Peterson for the SEM work.

REFERENCES

I - K.M. Prewo, J.J. Brennan, G.K. Layden, Am.Ceram.Soc.Bull. 65 (1986) 305-315, 320

2 - P.J. Lamicq, G.A. Bernhart, M.M. Dauchier, J.G. Mace, Am.Ceram.Soc.Bull. 65 (1986) 336-338

3 - M.W. Lindley, D.J. Godfrey, Nature 229 (1971) 192-193 4 - J.J. Brennan, Proc.BCRA Symp. Special Ceramics 6 (1974) 123-134 5 - J.D. Buckley, Am.Ceram.Soc.Bull. 67 (1988) 364-368 6 - B.E. Walker JR., R.W. Rice, P.F. Becher, B.E. Bender, W.Z. Coblenz,Am.Ceram.

Soc.Bull. 62 (1983) 916-923 7 - R.P. Boisvert, R.J. Diefendorf, Presented at 12th Ann.Conf. on Composites and

Adv.Ceram., Cocoa Beach, USA, Jan (1988) (to be published in Ceram. Eng. Sci. Proc.)

8 - D. Seyferth, Presented at ACS Int.Workshop Adv. in Si-based Polymer Science, Oahu, Hawaii, Nov (1987)

9 - G.E. Legrow, T.F. Lim, J. Lipowitz, R.S. Reaoch, Am.Ceram.Soc.Bull. 66 (1987) 363-367

10- F. Sirieix, Thesis University of Limoges, France 87-9 (1987) 11- F. Sirieix, P. Goursat, Rev.Int.Htes.Temp.et Refract., (1988) (to be published) 12- F. Sirieix, P. Goursat, A. Lecomte, A. Dauger, Composite Science and Technology

(1989) (to be published)

1200

01000 •

"'900 - - - - - -- - -- -- - -- - - --. - --- ------- ----

! ::J ... (Q650 .. Q) . Q.500

E 360 ------­Q) ~ 250 .--

15 20

Time ,h Figure 1. Temperature cycle for pyrolysis

I I

I

100

?R 90

Cl) ::s 80

"C .-en Cl) 70 -... TGA

.J::. 60 C) .-~ 50

100 200 300 400 500 600 700 800 900 1000

Temperature,OC

Figure 2. TGA curves for two different polymers: (0) Polymer treated in flowing nitrogen. (e) Untreated polymer.

Figure 3. SEM micrographs of the outer surface (a) and a fracture surface (b) of a pyrolysed composite (Bar = 10 J.'m)

97

98

Displacement Figure 4. Load/displacement curves for pyrolysed composites.

(A) Polymer with a high oligomer content. (B) Oligomer-free polymer and carbon coated fibres.

COMPOSITES A MATRICE CERAMIQUE, NOUVEAUX MATERIAUX A TRES HAUTES PERFORMANCES

J.C. CAVALIER, A. LACOMBE, J.M. ROUGES

Societe Europ8enne de Propulsion BP 37 ·33165 ST MEDARD EN JALLES • France

SEP has developed composite materials built with refractory reinforce­ments (carbon or ceramic) associated with a ceramic matrix. These thermostructural composite SiC/SiC and C/SiC materials combine a good resistance to oxidation and to chemical agressions with a high mechanical behaviour, at high temperature and an exceptional toughness­compared to the classical ceramics. News applications have been found for these new composites in the fields of engines on one hand and for very high temperature structures on the other hand. Very large industrial installations have been implemented since 1986 in order to meet the production requirements yielded by these new applications under development.

I-INTRODUCTION Forte d'une grande experience acquise dans Ie domaine des composites carbone-carbone depuis 1969, la SEP a developpe durant ces 10 dernie­res annees une nouvelle famille de materiaux composites dans lesquels sont associes un renfort fibreux carbone ou cerami que et une matrice ceramique. Ces materiaux ont Ie double avantage de presenter une resistance a l'oxydation nettement amelioree par rapport aux C/C et une tenacite tres superieure a celIe des ceramiques frittees. Nous nous proposons de montrer l'interet de cette famille, dans Ie cas particulier des materiaux a matrice SiC, en presentant : - leur mode d'elaboration - leurs principales caracteristiques - leurs applications et les pieces fabriquees ou en cours d'etude - les capacites industrielles mises en place a SEP.

2-FABRICATION DES MATERIAUX COMPOSITES A MATRICE SiC Dans Ie cas ou Ie renfort fibreux est un tissu, les operations de fabrication resumees sur la figure 1 comprennent les etapes suivantes: - decoupe des strates de tissu

100

- empilage de ces strates - mise en place des strates dans un outillage con9u specialement pour

les infiltrations gazeuses et permettant Ie contrale du taux de fi­bres desire : + environ 45 % avec Ie tissu de carbone + 40 % avec Ie tissu de carbure de silicium.

- consolidation de l'empilage de strates par un premier depot de la matrice SiC elaboree par un procede en phase vapeur

- demontage de l'outillage et poursuite de la densification par la repetition des cycles d'infiltration de SiC jusqu'a la densite souhaitee : + 2,1 pour Ie materiau 2D C-SiC + 2,5 pour Ie materiau 2D SiC/SiC.

La porosite ouverte residuelle est voisine de 10 ~~en volume pour les deux materiaux.

Figure 1 - Fabrication des composites a matrice SiC (C.V.I.) Parallelement aces materiaux adB5nfort bi-directionnel, la SEP a developpe les textures Novoltex a base de fibres de carbone. Ces textures multidirectionnelles sont caracterisees par un reseau de pores fins et reguliers , parfaitement adapte au procede de densifi­cation par infiltration en phase gazeuse. Elles presentent plusieurs autres avantages tels que : - une plus grande isotropie - une meilleure resistance au delaminage. Leur procede de fabrication, meca~isable et automatise, est totalement industriel. Les textures Novoltex peuvent etre elaborees en diverses epaisseurs faibles ou fortes. Elles sont utilisables pour realiser des pieces ayant de larges variations en epaisseur et des formes com­plexes monolithiques. Le taux de fibres objectif des materiaux est quasiment atteint lors de l'elaboration de la texture. Aussi leur densification ne necessite qu'un outill~ simplifie. Des textures multidirectionnelles SKINE~a tres hautes performances mecaniques sont actuellement developpees pour repondre specifiquement aux besoins de structures thermomecaniques formees de peaux et raidis­seurs integres.

3-CARACTERISATION DES MATERIAUX COMPOSITES A MATRICE SiC Les proprietes de ces 3 composites a matrice SiC sont donnees dans les tableaux des figures 2, 3 et 4.

Figure 2 - Caracteristiques du 2D - C-SiC Figure 3 - Caracteristiques du Novoltex C-SiC Figure 4 - Caracteristiques du 2D - SiC/SiC

La figure 5 montre l'evolution de la resistance en traction des trois composites en fonction de la temperature. Les mesures ont ete faites sous atmosphere neutre. II est interessant de constater qu'a 1400 o C, les resistances de traction sont superieures a 140 MPa. Ces differen­ces de comportement sont probablement dues aux evolutions des fibres a ces temperatures.

Figure 5 - Evolution de la resistance en traction en fonction de la temperature

La resistance a l'oxydation des composites a matrice SiC depend de nombreux facteurs tels que la temperature, la duree, et l'atmosphere.

101

Nous avons reporte sur la figure 6 la resistance en traction mesuree a temperature ambiante des materiaux siC/siC exposes durant vingt heures a l'air a differentes temperatures.

Figure 6 - Resistance en traction en fonction de l'allongement a rupture de SiC/SiC oxydes a differentes temperatu­res durant vingt heures

On observe que - la resistance en traction reste inchangee pour des temperatures de

vieillissement inferieures ou egales a 1200°C. Elle reste superieu­re a 140 MPa, meme apres 20 heures a 1550 oc.

- le module et le domaine elastique diminuent lorsque la temperature de vieillissement augmente de 800 a 1550°C. Par c~ntre, l'allonge­ment a la rupture est considerablement augmente de 0,3 a 0,8 %.

La figure 7 permet de mettre en evidence l'influence de la duree de vieillissement pour une temperature de 1200 0 C ainsi que celle d'une precontrainte de 130 MPa (valeur superieure a la limite elastique) avant vieillissement. On cons tate la encore que la resistance en traction reste inchangee apres vingt heures de vieillissement, avec et sans precontrainte initiale, et apres cent vingt heures sous air, alors que l'allonge­ment a rupture atteint une valeur proche de 1 %.

4-LES APPLICATIONS Par leur comportement thermomecanique et leur stabilite chimique, les CERASEP SiC/SiC et SEPCARBINOX C/SiC couvrent un large champ d'applications notamment dans les domaines de la propulsion a liqui­des, des moteurs aeronautiques et des structures thermiques d'avion spatial, principalement caracterises par: - des temperatures elevees - une ambiance oxydante.

4.1 MOTEURS DE PROPULSION BILIQUIDE

Les developpements ont porte sur les deux types de propulsion bili­quides, a ergols stockables et cryogeniques.

4.1.1 - Ergols stockables

Des chambres de combustion N204/MMH de diverses poussees (cf figu­re 8) (5 Na 6000N) ont ete testees, avec cyclages thermiques (jus­qu'a 400 cycles) et sur de longues durees cumulees (jusqu'a 36000s). Les principaux avantages apportes par le CERASEP siC/siC pour les durees les plus iongues {10000 a 100000 sl et le SEPCARBINOX C/SiC pour les durees moyennes (jusqu'a 1000 s) resident dans: - une temperature maximale de paroi de chambre admissible superieu­

re de 200°C a celle autorisee par les alliages metalliques gene­ralement utilises { 1550 0 C pour le siC/siC c~ntre l350°C (pour un alliage de Niobium) se traduisant par un gain d'impulsion specifique.

- une tenue en cyclage thermique fiable des composites a matrice SiC, par opposition a la fissuration possible des revetements de sili~uration obligatoire sur les alliages de Niobium pour les

102

proteger de l'oxydation, accroissant la surete de fonctionnement - des caracteristiques mecaniques specifiques tres elevees a hautes

temperatures, conduisant a un bilan masse favorable.

Materiaux SiC/SiC SiC/SiC SiC/SiC Novoltex C/SiC

Poussee (N) 5 5 200 6 000 Rapport melange 1,8 1,64 1,64 1,64 Pression (MPa) 1 1 1,6 1,2 Duree cumulee 36 000 24 000 1 000 870 Cyclage 400 Temperature paroi 1 700 1 700 1 600 1 500 interne (oC) Erosion au col 10 20 0 0

4.1.2 - Ergols cryogeniques

Dans Ie cadre du developpement d'un divergent composite pour mo­teur L02/LH2, SEP a conduit, au titre du marche Recherches et Technologies CNES, l'etude et la fabrication d'un divergent en SEPCARBINOX (Novoltex C/SiC) pour Ie moteur du 3eme etage d'ARIANE, HM7 (cf figure 9). Cette technologie, permet, outre Ie gain de masse par rapport au divergent metallique refroidi (~10 kg pour HM7) une augmentation de l'impulsion specifique. Des essais de divergents reduits en SEPCARBINOX ont mis en eviden­ce l'excellent comportement en oxydation de ce dernier, confortant l'objectif d'un essai nominal en simulation d'altitude du diver­gent HM7 echelle 1. Les conditions d'essais sont les suivantes :

Rapport de melange O/H 5,3 Temperature a la paroi 1 450°C Duree des essais 10 a 300 sec.

Divergent en SEPCARB C/C diminution de la masse de 4 % Divergent en SEPCARBINOX pas de variation de masse notable

4.2 MOTEURS AERONAUTIQUES

4.2.1 - Pieces de tuyeres

La SEP developpe en cooperation avec la SNECMA diverses pieces de reacteurs telles que volets de tuyere, cone d'echappement et an­neaux accroche-flammes (cf figures 10 - 11 - 12). Diverses experimentations sont menees sur ce type de pieces, dans des conditions reelles d'utilisation (essais sur moteur au banc). Les principaux resultats, au stade actuel du deroule des campagnes d'essais sont resumes dans Ie tableau ci-apres

103

Pieces Temperature (oC) Duree cumulee (h)

Anneau accroche-flamme 600/1150 12/1,7

Cone d'echappement 1000 320

Volets de tuyeres 750 245

Les points forts des CMC dans ce domaine d'applications resident dans : - un gain de masse obtenu par rapport a des solutions metalliques

(40%), consequence des caracteristiques specifiques elevees - une maintenance en service diminuee, en particulier pour les

pieces les plus sollicitees en chocs thermiques.

4.2.2 - Roues de turbines (figure 14)

Les roues de turbines sont soumises a un ensemble de sollicitations pour leRquelles les CERASEP SiC/SiC et SEPCARBINOX C/SiC sont bien adaptes. Des roues en SiC/SiC et C/SiC ont ete testees avec succes dans les conditions suivantes : - melange : air/kerosene - temperature : 1 150°C - duree : 1 H - vitesse de rotation: jusqu'a 70 000 t/mm L'amelioration de performance apportee pour l'utilisation des CMC concerne principalement I 'augmentation potentielle des temperatures de fonctionnement, et donc Ie rendement global du moteur, ainsi qu'un bilan masse favorable. Une caracterisation menee en parallele sur des eprouvettes repre­sentatives d'aubes a mis en evidence un bon comportement: - sous gradients thermiques (400°C entre bord d'attaque et bord de

fuite) - en choc thermique (passage de 400 a 1 200°C en quel~ues secondes) - en fatigue thermomecanique.

4.3 STRUCTURES POUR AVION SPATIAL

C'est en cooperation avec AMD-BA, sous contrat CNES, que la SEP deve­loppe des pieces de protection et de structures thermiques pour l'avion spatial HERMES.

4.3.1 - Le systeme de protection thermique

Le systeme de protection thermique (cf figures 15 - 16 et 17) est constitue en surface, d'un ensemble de panneaux en C/SiC ou SiC/ SiC assurant Ie profil aerodynamique de l'avion, dans les zones ou les flux thermiques sont eleves (intrados de voilure, flancs et dessus de la cabine).

104

L'isolation entre les panneaux et la structure froide etant rea­lisee par un complexe multi-couches de faible densite. Grace - a la masse optimisee des isolants internes et

- aux resistances specifiques elevees des C/SiC ou SiC/SiC, Ie bilan-masse d'une telle solution s'avere favorable par rapport au systeme des paves rigidifies tels que developpes sur l'Orbiter et qui par ailleurs restent limites en ~emperature.

4.3.2 - Les pieces de grandes dimensions pour structures thermiques (cf figure 16)

Compte-tenu de la petite taille d'Hermes, les faibles epaisseurs de profil en bout d'ailes ne permettent pas d'utiliser dans ces zones une structure froide, protegee par Ie systeme de protection thermique. L'architecture de l'avion spatial comprend donc differentes struc­tures thermiques telles que les derives latera1es, les drapeaux, les elevons et Ie volet ventral, soumises a des efforts gene raux importants (flexion, torsion et efforts combines) a hautes tempe­ratures. Les dimensions de ces pieces sont importantes, jusqu'a 2,50 x 1,80 m de surface. Le developpement des pieces de grandes dimensions a necessite la mise ~point par la SEP de textures specifiques quasi-3D, SKINE~ La demonstration de la faisabilite des PGD est basee sur la realisation, actuellement en cours, d'un caisson de derive en C/SiC de 1,8 x 0,8 x 0,3 m3. Cette piece de faisabilite, compati­ble avec la taille des installations industrielles existantes, sert de demonstrateur technologique des futures structures ther­miques d'Hermes. Le comportement de pieces representatives des applications pour l'avion spatial a ete evalue; en particulier un bord d'attaque SiC/SiC a ete teste avec succes jusqu'a 1400 0 C en ambiance oxy­dante, sous des sollicitations mecaniques de traction et compres­sion et des chocs thermiques representatif d'une mission Hermes la duree cumulee des essais avec une temperature superieure a 1 200°C est de l'ordre de 20 H. Une experimentation en cours sur une piece similaire en C/SiC de­vrait confirmer, a une echelle representative, l'aptitude de ce materiau a supporter des temperatures plus elevees (jusqu'a 1 700°C), demontree sur eprouvettes de caracterisation.

5-CONCLUSION SEP a developpe une famille de materiaux composites a matrice cerami­que a tres hautes performances thermomecaniques parfaitement bien adaptes pour realiser des pieces fortement sollicitees mecaniquement et fonctionnant a tres hautes temperatures en milieu oxydant. Des moyens industriels importants ont ete mis en place pour repondre aux productions de serie et aux realisations de pieces complexes et de grandes dimensions. En particulier, un four d'infi1tration de carbure de silicium de tres grande capacit~ (0 2,5 m x 2,5 m) est en cours de realisation pour Ie compte du CNES sous delegation de l'ESA.

Proprietes lunite! Temperature 23°C , 1 OOO°C ,

Taux de fibres % 45 Densite apparente 2,1 Porosite ouverte % 10

Resistance en traction MPa 350 350 Allongement a rupture % 0,9 0,9 Module de Young (traction) GPa 90 100

Resistance en flexion MPa 500 700

Resistance en compression II MPa 580 600

! MPa 420 450

Resistance en cisail. interlam. MPa 35 35

I -6 ,

Diffusivite thermique II 10 nM 11 7

1 I 5 2

I I _, -1'

Coef. de dilatation II 110 K I 3 1 I I 5

I I , I

"Tenacite" K1R I MPaVrnI 32 32 I I

I

Figure 2 - Caracteristiques du 2D C-SiC

Tissu

Oecoupe de strates

Montage dans l'outillage

Obtent.ion d'une preforme consolidee Pre forme finie densifiee

'----~ ~ ~

Figure 1 - Fabrication des composites a matrice SiC (C.V.!.)

105

1 400°C

330

100

700

700 500

35

8 2

I I I I r

32 I I

106

Temperature Proprietes 23°C I 1 OOooC I 1 400°C

Taux de fibres Densite apparente Porosite ouverte

Resistance en traction Allongement a rupture Module de Young ( traction

Resistance en flexion

Resistance en compression II " " J..

Resistance au cisail. inter lam

Diffusivite thermique /1 " "J.

Coeff. de dilatation II " " ...L

% 24

GPa 75 85

MPa 300 -MPa 650 I 700 MPa 740 740

MPa 100 40

1~~1! 12 5 I " I 9 4

11o'lI!{..t I 2 2,5

Figure 3 - Caracteristiques du Novoltex C-SiC

I I

Proprietes Temperature 23°C I 1 OOooC I

Taux de fibres % 40 Densite apparente Porosite ouverte

Resistance a traction Allongement a rupture 0,3 0,4 Module de Young (traction) GPa 230 200

Resistance en flexion MPa 300 400

Resistance en compression II MPa 580 480 " " .l.. MPa 420 I 380

Resistance au cisail. interlam MPa 40 35

Diffusivite thermique /,1 lo'~; 12 I 5 " " I " I 6 I 2 I

Coeff. de dilatation I I lO"'o~1 3 " " ..L " 2,5

"Tenacite" K1R MPa~ 30 I 30

Figure 4 - Caracteristiques du 2D SiC/SiC

140 0,8 70

-800 770

40

5 4

-

1 400°C

150 0,5

170

280

300 250

25

5 2

-

30

I I I

107

i .-'''' I i,) (' .~ (. ----, I ............ I,oI • • t -:;1

"" ..... ,- ' ....

Figure 5 - Evolution de la resistance en traction en fonction de la temperature

200

180

so

u.s . TRACTION 1M,., OIl'ICTlOM I AMCI 2

.''200 ·c

"" so 90

Figure 6 - Resistance en traction en fonction de l'allongement a rupture de SiC/SiC oxydes a differentes temperatures durant 20 H

RES. TRACTION

IMI'aI OIl'ICTIOM I AlII D 2 200

\

180 :;; t:.":,~:tIi ~

120 /;> ao ~" '~,v-..,

IU.L • RUPTURE I~

.10 . .31 ..30 .40 .!It .IG .10 .111 .90 UlO

Figure 7 - Resistance en traction en fonction de l'allongement a rupture de SiC/SiC oxydes a 1200°C

108

Figure 8 - Chambres de propulsion biliquides

Figure 9 - Divergent HM7 en Sepcarbinox

Figure 10 - Volets de tuyere

109

Figure 11 - Cone d'echappement

Figure 12 - Anneau accroche-flamme

Figure 14 - Roue de turbine

110

PANNEAUX

Figure 15 - Panneaux

STRUCTURES THERMIOUES

Figure 16 - Structures thermiques pieces de grandes dimensions

Figure 17 - Bords d'attaque

RESISTANCE TO CRACK GROWTH IN FIBRE REINFORCED CEMENT:

ABSTRACT

EFFECT OF FIBRE PROPERTIES

L. DESCHRYVER, A.A. BUNSELL, A. LE FLOCH'

Ecole des Mines de Paris Centre des Materiaux - BP 87 - 91003 EVRY Cedex - France

'Everite - BP 11 - 77190 DAMMARIE-LES-LYS - France

Slow stable crack growth is a prominent feature of the fracture behaviour of fibre reinforced cement. The main mechanism for resistance crack growth during crack extension in this type of composite is microcracking ahead of the crack tip. This phenomenon is characteristic of a cement matrix, containing many pre-existing defects which will propagate during loading. Its effects can be calculated by applying linear elastic fracture mechanics to the microcracks. Fibres stabilise the crack propagation as long as they bridge the crack. A statistical study, based on fibre bundle behaviour, allows an explaination of the effect of fibre properties on crack growth resistance in this type of composite.

1. INTRODUCTION

Abestos-cement is a widely used material in the construction industry. It is made up of a small percentage of short asbestos fibres randomly distributed in a cement matrix. Due to health hazards ascribed to asbestos, there is a need for suitable substitute fibres for cement reinforcement. A large number of fibres, with very different mechanical properties, have been proposed. This diversity underlines the uncertainties still existing concerning the reinforcement mechanisms of cement matrices.

The role-played by short fibres in a brittle matrix such as cement is in the control of crack propagation. Therefore the increase of the crack growth resistance during crack extension is a good indication of reinforcement effects.

112

The aim of this study was to understand the influence of fibre properties on the composite fracture resistance. Three types of synthetic fibres out of the most promising candidates were selected each with very different mechanical properties: glass, polyvinyl alcohol and acrylic.

As a first step, the mechanical behaviours of these fibres, by a single fibre tensile test were studied.

Then, the crack growth resistance of associated composites was measured. All specimens were cut from flat sheets made on a pilot plant Hatschek machine, in conditions similar to those used in the industrial process.

This process does not allow manufacture of sheets consisting only of cement and synthetic fibres, but requires the addition of a "process fibre", such as cellulose or asbestos.

The reinforcement effects of cellulose and asbestos fibres alone have also been studied. The compositions of materials tested are given in Table 1.

2. THEORY

Synthetic fibres

Fibres, like other materials, contain many defects, which influence their mechanical properties. The results of tensile tests, can be very dispersed, and it is necessary to determine their statistical distribution.

Testing the fibres at different gauge lengths is a way to select the best statistical model describing their behaviour. If an effect of specimen length exists, the shortest being the strongest, the hypotheses of the weakest link theory are verified and the Weibull statistical distribution will be a good mathematical representation. If not, we shall demonstrate that a Gaussian distribution can be used.

Composites

In fibre-cement composites, the matrix always fails before the fibres. Fibres cannot significantly delay the initiation of matrix damage. Their effect appears essentially during matrix crack extension. So, the increase of crack growth resistance with crack length is characteristic of the fibre reinforcement of cement.

This increase is generally explained by two mechanisms - micro cracking ahead of the crack tip, - fibres bridging the crack and limiting its opening.

Microcracking is usually studied by analogy with the plastic deformation of metals. Microcracking can also be described as a function of the energy dissipated during the creation of new surfaces,

113

such as in macro-cracks. Micro and macro-cracking can therefore be studied by using linear elastic fracture mechanics.

Microcracking in cement matrices is a consequence of microstructural defects, like pores or microcracks due to shrinkage. The size of these defects is about 1 mm, much larger than fibre radii. Fibres interact with microcracks, and stabilize their propagation. The mechanics are exactly the same with a macrocrack.

The study of the behaviour of these microcracks under a known stress needs some assumptions :

- linear elastic fracture mechanics can be applied. - there is no interaction between microcracks. - microcracks are stressed essentially in a tensile mode.

A microcrack ahead the macrocrack tip can therefore be considered as being similar to a single crack in an infinite sheet, approximated by a eN specimen.

The propagation of a crack is supposed stable so long as fibres cross it

- if a fibre bundle across the microcrack is broken, the stress

intensity factor will be : K = uapp fma. where Uappis the applied stress and 2a the size of the microcrack.

- if not, the fibres will support the whole stress, and this factor will be : K = o.

Knowing the fracture probability Pf of the fibre bundle, we can calculate the average stress intensity factor for a microcrack of a given length under a given stress.

This probability is linked to the fracture probability of a single fibre. The problem of fibre bundles is well known, and can be resolved for bundles made of a great number of fibres: the maximum load supported by a bundle is attained for a given percentage of broken fibres, depending on the statistical distribution of single fibre strengths (Fig. 1).

The present case considered is somewhat different. The number of fibres crossing a microcrack is small. If we assume that the defects are circular, it is between ten and fifty. So, the stress concentration on fibres adjacent to a broken fibre cannot be neglected.

The number of adjacent fibres is independant of the size of the defect. After the fracture of a fibre, the stress concentration on adjacent fibres will not depend on defect size.

So, the maximum percentage of broken fibres.

load on fibres,

these bundles is not reached for a but for a critical number of broken

114

3. EXPERIMENTAL WORK

Synthetic fibres

All tests have been performed on a specially designed single fibre tensile test machine. The three different gauge lengths choosen - 20 mm, 10 mm, and 5 mm - are as close as possible to the lengths used for cement reinforcement. This procedure allows the statistical distribution of fracture strengths to be determined.

A characteristic of the Hatschek process used to produce the samples is a strong tendency for self-alignement of fibres in the machine direction. This effect is very pronounced with synthetic fibres, which are longer than natural ones.The mechanical properties of composites become significantly different in the machine direction from those in the transverse direction. To obtain more isotropic materials, we manufactured sheets the thicknesses of which were a quarter of the final thicknesses, and stacked them with different orientations. Composites obtained were orthotropic and stratified in a L-T-T-L sequence.

The compact tension (CT) geometry of specimens used in this study is illustrated in Fig. 2. In order to guide the fracture, side grooves were machined thereby reducing the thickness by 40-50 %. Due to these grooves, we were able to oberve the propagation of cracks perpendicu­larly or parallel to the fibres in a globaly nearly isotropic material.

Crack growth resistance was calculated employing linear elastic fracture mechanics. Its increase with crack extension is represented by R-curves.

We have also monitored acoustic emission from these CT specimens during loading. The use of two transducers permitted the location of the source of emission. In this way, we could measure the size of the damaged zone ahead the crack tip during crack extension.

4. EXPERIMENTAL RESULTS

The glass fibres showed brittle behaviour and the gauge length had a significative effect on their tensile failure stresses. Weibull statistic seems to be a good means of describing the dispersion of results. The Weibull coefficient was found to be nearly 3.5.

The organic fibres behave in a ductile manner and failure strains were greater than with glass fibres.The gauge length did not change the average failure stresses. Using all the results, we verified that their distribution was Gaussian.

Composites

All materials studied showed stable crack growth, and resistance increased with crack length.

115

Matrix degradation began prior to macrocrack extension : we could detect some acoustic emission ahead of the macrocrack tip before its extension.

The size of the damaged zone increased during the crack propaga­tion, to reach a maximum stable value. All cellulose-cement and synthe­tic fibre-cellulose-cement studied gave the same results : 20 mm at the crack initiation, and 60 mm as the maximum value. The only difference observed was for abestos-cement : the size of the damage zone was smaller : 12 mm at the begining and 30 mm at the end of the test.

Fibre orientation had no effect on the size of the damage zone, and no effect on resistance to crack initiation. Differences appeared during the propagation.

When the crack was parallel to the fibres, resistance quickly reached a threshold, whilst if perpendicular, crack resistance continued increasing as is shown in Figure 3. Considering the size of the damage zone, we believe that the maximum value measured with this fibre orientation was not far from the threshold value.

Amongst the synthetic fibres, best results are obtained for polyvinyl alcohol fibres. Glass fibres, the failure stress and Young's modulus of which are greater, give lower strength composites. The worst results were obtained with acrylic fibres which had failure strains greater than PVA or glass fibres (Table 2).

Observations of the fracture areas showed some evidence of fibre fracture for each synthetic fibre.

1 With natural fibres, previous observations were confirmed that : asbestos-cement does not resist crack growth very well and is worse

than cellulose-cement, despite the fact that asbestos fibres are stronger than cellulose fibres (Table 2).

5. DISCUSSION

The increase of resistance to crack growth, for a crack running parallel to the fibres, is exclusively an effect of the microcracking ahead of the crack tip.

Synthetic fibres are unable to enlarge the damaged zone, or to delay the beginning of propagation of the preexisting defects. But they make propagation more difficult.

They act in the form of small bundles crossing defects as described in section 2. Examining the behaviour of these bundles, with a critical number of fibres broken rather than a critical percentage, we can calculate the composite crack growth resistance in cpmparison with the cellulose-cement matrix.

116

Therefore the single fibre tensile strength is the parameter which has the most influence on the resistance to crack growth in this type of composite. The effect is a function not only of the average fibre strength, but the whole statistical strength distribution of the fibres in the bundle.

This explains why PVA fibres give better results than glass fibres as the latter shows much wider scatter in their strength.

Considering bundles in the way proposed gives a non-linear relation between the composite resistance to crack propagation and the volume of fibres. This result has been observed experimentally for asbestos and cellulose fibres.

The smaller damage zone in asbestos-cement is a consequence of the fineness and high modulus of asbestos fibres. This high modulus delays loading of the cement matrix, and damage initiation and the fineness reduces the average distance between fibres.

The most important damage in abestos-cement is fibre pull-out as the fibre length does not obtain the critical length.

Fibre pull-out can be studied like fibre fracture. The dispersion of values is negligible compared to structural dispersion, due to the relative positions of fibre and crack. Fibre pull-out has a maximum dispersion ; bundles are always' weaker than the strength of fibre bundles, for the same average stress.

This disadvantage in the case of asbestos is made worse by the fineness of asbestos fibres, increasing the number of fibres for each microcrack.

BIBLIOGRAPHY

1. R.N. Swamy and H. Starrides, "Influence of fibre reinforcement on restrained shrinkage and cracking". ACI Journal, 76, 3, (1979), 443. 2. Y.W. Mai, M.l. Hakeem and B. Cotterell, "Effects of water and bleaching on the mechanical properties of cellulose fibre cements", J. Mat. Sci., 18, (1983), 2156. 3. A.J. Majumdar and V. Laws, "Fibre cement composites: research at BRE" , Composites (Jan. 1979), 17. 4. J. Aveston and A. Kelly, "Theory of multiple fracture of fibrous composites", J. Mat. Sci. 8, (1973), 352. 5. D.J. Hannant, D.C. Hughes and A. Kelly, "Toughening of cement and other brittle solids with fibres". Phil. Trans. R. Soc. A310. 6. J.P. Romualdi, "The strengthening of brittle materials". Mat. Sci., and Eng., 15, (1974), 31. 7. J.C. Lenain and A.R. Bunsell, "The resistance to crack growth of asbestos cement". J. Mat. Sci., 14, (1979). 321.

117

0 1 2 3 C

/ 6.0 % 8.1 % 0.6 % Cellulose A

/ 6.2 % 8.9 % 10.9 % Asbestos P o % 2.1 % 2.8 % 3.4 % Polyisnyalcohol

= 6 % cellulose

G o % 1.9 % 2.6 % 3.1 % Glass

Ac o % 2.3 % 3.1 % 3.8 % Acrylic

Table 1 Fibre composition of specimens

Transverse Direction Machine Direction

KR MPa~ KR MPa~

P 0 2.3 3.5 P 1 4.0 7.3 P 2 4.3 8.2 P 3 6.1 9.5

AC 0 1.9 3.4 AC 1 2.2 5.6 AC 2 2.3 6.0 AC 3 2.6 6.4 G 0 2.6 3.8 G 1 3.6 6.0 G 2 3.4 6.3 G 3 4.0 7.1 C 1 2.2 3.4 C 2 2.4 4.7 C 3 3.2 5.6

AS 1 1.9 2.4 AS 2 2.6 3.3 AS 3 2.5 3.9

Table 2 Crack resistance of the composite specimens, as defined in Table 1.

118

% broken fibres

1

. 5

o

Figure 1

Pmax Damage accumulation in a fibre bundle as a function of load applied to the bundle and catastrophic failure.

6

160

~ 1p

.. tP I I

I

0

a/w = 0.3

BIB1 = 0.6

HI. = 0.4

1 H

W

B

II U I I

B1

• = 125 mm Bl = 5 mm

119

100

Figure 2 : Geometry and relative dimensions of the specimens employed.

120

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THERMO·MECHANICAL CHARACTERIZATION OF CERAMIC COMPOSITES MADE OF

A LAS GLASS·CERAMIC MATRIX REINFORCED WITH SILICON CARBIDE (NICALON) FIBERS

ABSTRACT:

E. MENESSIER, A. GUETTE, R. PAILLER, R. NASLAIN, l. RABARDEL, B. HOSTEN', T. MACKE", P. LESPADE'"

Laboratoire de Chimie du So/ide du CNRS , Laboratoire de Mecanique Physique

•• Laboratoire de Genie Mecanique (fUT-A) 351 cours de /a Liberation - 33405 TALENCE - France

••• Aerospatia/e Aquitaine BP 11 - 33165 ST MEDARD EN JALLES - France

Pseudo-unidirectional composites made of SiC Nicalon fibers and LAS glass-ceramic matrix were obtained via a low temperature sol­gellhot pressing route. The fibers have been coated with a thin layer of pyrocarbon in order to promote fiber pull out. The elastic constant Cij matrix and engineering elastic moduli at room temperature were derived from US wave propagation experiments. The failure strength was measured for different processing conditions through three point bending testing. The failure energy and dynamic toughness were calculated from instrumented CHARPY test data. Thermal expansion experiments were performed on both unreinforced LAS matrix and SiC / LAS composites.

KEYWORDS:

CERAMIC MATRIX COMPOSITES ELASTIC CONSTANTS LITHIUM ALUMINO SILICATE THERMAL EXPANSION

1· INTRODUCTION:

GLASS CERAMICS TOUGHNESS SIC NICALON FIBER

Ceramic matrix composites (CMC) made of silica-based glass or glass-ceramic matrices reinforced with either continuous or short ceramic fibers (carbon or silicon carbide) are a subject of more and more research with a view to applications at medium temperatures (i.e. up to BOO-IOOO°C). With respect to the corresponding unreinforced matrices, such composites exhibit improved stiffness, failure strength and toughness up to temperatures close to the Tg transition of the matrices.

122

Futhermore, they show dimension stability still better than those of the unreinforced matrices due to the low coefficients of thermal expansion (CTE) of both the fibers and matrices (particularly when carbon fibers, which have negative CTEs are used).It has been established very recently that, in such brittle-brittle composites, the fibers should not be too strongly bound to the matrix in order to achieve a high toughness. Such a requirement is fulfilled when a thin layer of low shear strength material (e.g. pyrocarbon) is present at the fiber-matrix interface and when the matrix CTE is lower than the radial CTE of the fiber. In silicon carbide fiber composites, this thin film of pyrocarbon is formed in-situ as the result of some fiber-matrix interaction during processing, if temperature is high enough [1].

The aim of the present contribution is to report the results of mechanical testing and physical characterization which have been performed on pseudo-unidirectional composites made of SiC (Nicalon) fibers embedded in a LAS glass-ceramic matrix, according to a low temperature sol-gel/hot pressing route. The study has been focussed on the effect of some processing parameters (e.g. the thickness of a pyrocarbon interphase deposited on the fibers prior to impregnation) on the mechanical characteristics of the composites (elastic constants, failure strength, dynamic toughness).

2 - EXPERIMENrAL:

The materials used in this study are pseudo-unidirectional composites made of a LAS glass-ceramic matrix (with a composition close to that of 13-spodumene Li20 A1203 4Si02) reinforced with SiC based fibers (mainly Nicalon NLM 202 from Nippon Carbon).They were obtained, at medium temperatures (i.e. less than 900°C), according to a sol-gel route which has been described in details elsewhere [2]. The starting material was a pseudo-unidirectional fabric made of Nicalon fibers running in the warp direction (referred to as the 1 direction) and maintained parallel one another with a small amount of fibers (6%) running in the woof direction (referred to as the 2 direction). The processing technique involved three main steps : (i) the fabrics were first impregnated with a sol and dried, the sequence impregnation / drying being repeated several times in order to achieve an overall fiber volume fraction of = 50%, (ii) then, the prepreg sheets were stacked together with the same fiber orientation and hot pressed at about BOO°C and (iii) the matrix was finally ceramed.

Since the aim of the present study was mainly to identify the effect of some processing parameters (e.g. the thickness of the pyrocarbon interphase ) on the mechanical or thermal behaviors of the composites, the tests were limited to the measurements of: (i) the elastic constants by US wave propagation, (ii) the failure strength in 3 points bending (iii) dynamic toughness by CHARPY impact testing and (iiii) thermal expansion.

123

3 - RESULTS AND DISCUSSION:

3.1.Elastic constants and enltineerin~ elastic moduli.

The engineering elastic moduli (longitudinal and transverse Young moduli ,shear moduli )of the composite (considered as a transverse isotropic material ) were calculated from the values of elastic constants derived from US wave propagtion experiments on the basis of theoritical models reported by B.Hosten and al. [3]. Their variations as a function of the carbon interphase thickness are represented in fig.1 . There is a significant decrease in the values of E1 ,E2 , G23 and G31 when the thickness of the carbon interphase increases from 0 to 0.4 ~m [4]. These results show that the expected effect, i.e. a softening of the material with respect to stresses resulting in mode II failure mechanisms (stress relaxation at crack tip, deflexion of mode I crack propagation along the fiber-matrix interface), has been achieved.

3 .2. Mechanical behavior in three point bendin~ loadin~.

The effect of some processing parameters such as the degree of ceramization of the matrix and the nature of the fiber-matrix interface (presence or not of a pyrocarbon interphase of a given thickness), on the mechanical behavior at failure was studied in three point bending at room temperature (experimental conditions:spanlthickness ratio of 25,and deflection rate of 0.1 mm/mn).

As shown in fig. 2, the mechanical behavior of the composites is typically brittle (i.e., low failure strength and work of fracture, no fiber pull out) when the LAS matrix has not been ceramed after densification. These features suggest that the fibers are strongly bonded to the matrix and are in agreement with the results reported for related systems. However, it clearly appears that such a behavior is no longer observed when the LAS matrix has been ceramed .

As a matter of fact, an increase in the composite failure strength occurs and fiber pull out takes place. Since the ceramization temperature seems to be too low to result in the in-situ formation of a pyrocarbon interphase at the fiber-matrix interface, the observed improvement in the mechanical behavior of the composites could be simply related to the numerous grain boundaries and other microstuctural defects due to the ceramization of the vitreous matrix.Such defects interact with propagating cracks and modify the failure mechanisms [5] . The values of the failure strength measured here on SiC (Nicalon) / LAS composites with no treatment of the fiber­matrix interface are low with respect to those reported by other investigators for materials obtained at higher temperatures from slurries and more complex matrices [6]. They could be explained by :(i) a too strong fiber-matrix bonding and (ii) some fiber degradation during processing (some diffusion of Li and AI within the fiber seems to occur, even at the low temperatures which have been used here , on the basis of SIMS analyses) as already reported for similar composites [1,7].

124

Three point bending tests were then performed on pseudo 1D-SiC (Nicalon) / pyrocarbon interphase / ceramed LAS composites.The values of flexural strengthes are given in fig.3 for three different pyrocarbon thicknesses. The typical load / midspan deflexion curve is shown in fig.2 for a composite with a pyrocarbon interphase 0,4 11m thick. As expected, the mechanical behavior is less brittle (with respect to that of the composite with no interphase). A significant increase in failure strength and toughness is observed with a lot of fiber pull out (fig.4), in agreement with the results previously reported by different authors for similar materials processed according to the high temperature route [1,6,S].

Thus the low temperature sol-gel route can be used for the processing of SiC / LAS composites exhibiting a non brittle behavior. However, such a behavior is achieved only when the SiC fibers have been coated with a thin layer of pyrocarbon prior to the impregnation step. In addition, to its mechanical role (i.e. crack blunting and crack deflexion) the pyrocarbon probably also acts as a diffusion barrier (with respect to Al and Li).

3.3. Dynamic toughness from instrumented CHARPY tests.

Dynamic toughness measurements have been performed on a pseudo 1D-SiC / Pyr. C / ceramed LAS composite in which the thickness of the pyrocarbon interphase was 0.4 11m with an instrumented CHARPY equipment according to a technique which has been described elsewhere [9]. The failure energy UD as well as the dynamic toughness (KID and KIDM) were derived from the load-time curve shown in fig.5. Assuming that such composites can be described as elastic damageable materials, the R curve was derived from the load-time curve [9]. As shown in fig.5 the energy which is necessary for crack propagation increases when the length of the crack is raised (due to the effect of numerous micro­damaging mechanisms) a feature which is commonly observed in most CMC with fibers weakly bonded to the matrix.The values of Un and KIDM

of this material are 12.S KJ/m2 and 22 MPa/m l/2 respectively. It appears that SiCILAS composites processed according to the sol-gel low tempe­rature route have a toughness similar to that of the composites obtained via the slurry high temperature route, provided a pyrocarbon interphase has been deposited on the fibers prior to the impregnation step [6,10,11].

3.4. Thermal expansion.

An important advantage of the LAS based materials lies in their low CTE. Moreover, since a LAS matrix can be regarded as an isotropic medium, the radial (arr) and axial (aaf) CTE of the Nicalon fiber (which are important parameters in the modeling of the thermomechanical behavior of CMC) could be derived from thermal expansion data measured on the pseudo 1D-SiC / LAS composites. The thermal expansion of an unreinforced LAS specimen and a p-lD-SiC / LAS composite within the 20 - 550°C temperature range (fig.6), were performed under a nitrogen

125

atmosphere. The values of CTE obtained were 25x10-7o Col for the LAS matrix, 26x10-7 °C-I and 25x10-7 °C-I for the composite in the 1 and 2 directions respectively.Assuming the fibers bonded to the matrix through a strong enough bonding ,the axial and radial CTE of the Nicalon fiber were calculated from the matrix and composite constants (the elastic moduli and POISSON's ratios obtained from US wave propagation measurements and CTE ) by applying Chamberlin's or Shapery's models [12,13]

The values of (laf (i.e. 28x10-7°C-1) is close to the value given by th

producer (i.e. 31x10-7oC-1) and arfis about 25x10-7 °C-I . It comes out that, as expected from the crystal structure of B SiC and processing technique, the Nicalon fiber is almost isotropic within the 20 - 500°C temperature range, as far as thermal expansion is concerned and the above models applicable.

4 - CONCLUSION

The present study has shown that tough SiC / LAS composites can be obtained according to a sol-gel low temperature route if the fibers have received a coating of a soft material (e.g. pyrocarbon) prior to the sol impregnation step. The pyrocarbon protects the brittle fibers against the notch effect due to matrix microcracking, increases the work of fracture and to a less extend plays a role of diffusion barrier. Better results, from the standpoint of mechanical behavior and interaction with the environment, could still be obtained by an optimization of the fiber coating.

ACKNOWLEDGEMENT This work has been supported by AEROSPATIALE, etablissement d'Aquitaine whose authorization for publishing this study is acknowledged.

REFERENCFS

1- Brennan J.J.,"Tailoring Multiphase and Composite Ceramics", (R E. TESSLER et aI, eds.), Material Science Research,20 (1986) 549-560, Plenum press New-York.

2- Menessier E., Guette A.,Pailler R, Naslain Rand Lespade P., Brevet fran¢s nO 87-18023

3- Hosten B.,Deschamps M. and Tittmann B.R. ,J. Acoust. Soc. Am.,8215 (1987) 1763-1770

4- Menessier E.,Guette A.,Pailler R.,Naslain R.,Hosten B. and Lespade P. , Proc. JNC 6 - Paris 11-13 Oct. 1988 (J.P. FAVRE and D.VALENTIN eds.) pp.195-210

5- Rice RW. ,Ceram. Engineer. and Science Proc., 6 (1985) 589-607 6- BrennanJ.J.and Prewo KM. ,J. of Mater. Sc.,17 (1982) 2371-2383 7- Menessier E., Guette A. , Pailler R. and Naslain R,(Submitted to

Ceram. Eng. and Sc. Proc.)

126

8- Chaim R. and Heuer A. H. ,Advanced Ceramic Materials, 2/2 (1987) 154-15

9- Macke T., Balette J.J. and Quenisset J.M. , to be published in Proc. Impact 87, Deutsche Gesellschaft fur Metallkunde EV,1987,Oberursel .

10- Mah T.,Mendirata M.G.,Katz A.P. ,Ruh R. and Mazdiyasni K.S.,J. of Amer. Ceram. Soc., 68/1 (1985) C27-C30

11- Marshall D.B. and Evans A.G. ,J. of Amer. Cera. Soc., 68/5 (1985) 225-231

g .., .. .2

12- Shapery R.A. ,J. of Composite Materials, 2 (1968) 380-391 13- Chamberlain N.J. , BAC report, SON (P) 33, November 1968.

e (pm)

-0- E2

200 ...... E1

--- G23 ...... 31

0,1 0,2 0,3 0,4 0,5

60

SO

CO

30

20

0 .6 0 .8 1.0 1.2 1._ 1.6

Figure 1: Variations of the elastic moduli as a fonction of the carbon interphase thickness .

Figure 2: Load-deflexion curves of:

a) p-1D SiC/ vitreous LAS

b) p-1D SiC/ ceramed LAS

c) p-1D SiC/ C/ ceramed LAS composite.

127

e (f.I m) 500

400 III

'"""' 000 co P-t

III ~ ax:l '-'

III

100

O~~~~~~~~~W

0,0 0,2 0,4 0,6 0,8 1,0 1,2

Figure 3: Variations of the fracture strength in three points bending loading of a p-1D SiC/CILAS composite with the thickness of the pyrocarbone interphase.

Figure 4: Failure surface of a p-1D SiC / C / ceramed LAS composite.

'~.---------------------,----------------. " lime hniC'rosccondes) crncK h.'u;;lh (1I1In)

"

Load-time curve R curve " .. 3

~ . DO

' ..

C o ".

Figure 5: Instrumented C?arpy testing of a p-1D SIC / C / LAS composite.

Figure 6: Thermal expansion curves of a p-1D SiC / LAS composite .

"

~T---------------------~

100

o~~~~~~-L~~--~~~

o 100 :m 3Xl 400 500 00)

5 <I

INTERFACE CHARACTERISATION BY TRANSMISSION ELECTRON MICROSCOPY AND AUGER ELECTRON

SPECTROSCOPY IN TOUGH SIC FIBER (NICALON)·SIC MATRIX COMPOSITE WITH A BORON NITRIDE INTERPHASE

O. DUGNE, S. PROUHET, A. GUETTE, R. NASLAIN, J. SEVEL Y·

Laboratoire des Composites Thermostructuraux Europarc - 3 avenue Leonard de Vinci - 33600 PESSAC - France

*Laboratoire d'Optique Electronique 29 rue Jeanne-Marvig - BP 4347 - 31055 TOULOUSE Cedex - France

ABSTRACT

The toughness of SiC fiber (Nicalon) /SiC matrix composites is increased when an interphase, made of a soft material (e.g. pyrocarbon or boron nitride) and playing the role of a mechanical fuse, is present at the fiber-matrix interfaces. The BN-interphase, deposited from a gaseous BF3-NH3 precursor, has been analyzed both chemically and microstructurally by TEM and AES, as well as the associated fiber­interphase and matrix-interphase interfaces. The sequence of materials observed at the fiber-matrix boundary is SiC matrix/BN interphase/Si02/carbon/Nicalon fiber.

1- INTRODUCTION

Ceramic matrix composites (CMC) made from SiC-based Nicalon fibers often exhibit a brittle mechanical behavior unless the fiber-matrix interfaces have been properly optimized during the high temperature step of their processing. It is now well accepted that tough CMC are obtained when a thin layer of a soft material is formed at the fiber­matrix interfaces as the result of : (i) some chemical reaction taking place in-situ at high temperature during processing or (ii) a treatment of the surface of the fibers (e.g. a CVD coating) prior they are embedded in the matrix. Among the soft materials which have been suggested as an interphase in SiC-, Zr02-, mullite- or glass-ceramic matrix composites, carbon and boron nitride, both having a layered structure, are those which result in the best mechanical properties in terms of strength and toughness /1-9/. However, the detailed mechanism according to which the interphase mechanical fuse works still remains imperfectly understood.

KEY WORDS : SiC/SiC COMPOSITES, BN INTERPHASE, INTERFACE, TEM, AES

130

The aim of the present contribution is to identify, on the basis of point analyses at a submicron scale (mainly TEM and AES), the sequence, composition and microstructure of the materials which are present at the fiber-matrix boundary in SiC (Nicalon) -SiC CMC made according to the chemical vapor infiltration (CVI) process.

2- EXPERIMENTAL

The materials which have been studied were samples of SiC-SiC composites prepared from a SiC-based fiber (Nicalon) preform according to a two step CVI procedure. In a first step, the preform was coated by a thin layer (i.e. less than 1 pm) of pyrolytic hexagonal boron nitride deposited from a BF3-NH3 gaseous precursor according to a technique which has been described in details elsewhere /10, 11/. In a second step, the BN-coated preform was densified by SiC deposited from a CH3-Si-C13 (MTS)-H2 gaseous precursor, according to the procedure reported in /12-14/. The microstructure of the samples was studied by transmission electron microscopy (TEM)(*) on thin foils which have been obtained by mechanical grinding and ion milling (**). The chemical analysis were performed both by Electron *:~ergy loss spectroscopy (EELS) and Auger electron spectroscopy (AES) .

3- RESULTS AND DISCUSSION

3.1- The pyrolytic boron nitride interphase The thin pyrolytic BN interphase was characterized by TEM in a SiC

(Nicalon)/SiC (CVI) composite from the standpoints of microstructure and composition (fig. 1). The data are given in table I and are compared with those obtained on BN-coated Nicalon fibers which have been extracted from a preform prior to the SiC-CVI densification /15/. It appears that the SiC-isothermal CVI infiltration acts as an annealing treatment regarding its long duration (i.e. several hundreds of hours) and results in : (i) a decrease in the interlayer d002 distance, (ii) a lowering of the oxygen content of the £ilm and (iii) a N:B atomic ratio which tends towards the BN stoichiometry. Features (ii) and (iii) seem to be related to a simultaneous evolution of oxygen and boron probably as gaseous boron oxides. On the contrary, the SiC­CVI isothermal densification does not modify the turbostratic character of the pyrolytic BN, the BN layers remaining randomly orientated in the core of the film and almost aligned parallely to the fiber surface through a thickness of about 10 nm (fig. 1 and 2).

3.2- SiC-matrix/BN and BN/Nicalon fiber interfaces As shown in fig. 1, there is no interaction zone between the BN

interphase and the SiC-CVI matrix, a feature suggesting that no chemical reaction takes place between the BN surface and the MTS-H2 precursor at the infiltration temperature (i.e. at about 1 OOO·C).

On the contrary, the BN-Nicalon fiber interface appears much more complex. As a matter of fact, two thin sublayers are present at the fiber-interphase boundary : (i) the first, which is in contact with the fiber itself, has a thickness of about 80 nm and a white appearance

(.) Philips EM 400 T and Jeol 200 CX ; ( •• ) Gatan 600 B (***) PHI 590 AES microprobe

131

while (ii) the second, in contact with BN, has a thickness of about 40 run and a black appearance (fig. 1). On the basis of the EELS analyses and AES depth profiling, the two sublayers were respectively identified with pyrocarbon and almost pure silica. Moreover, electron microdiffraction experiments have shown that both of them were amorphous. Therefore, the following sequence of materials are present at the fiber-matrix boundary SiC-CVI matrix/pyrolytic BN interphase/vitreous silica/amorphous carbon/Nicalon fiber. It is, noteworthy that the amorphous silica and carbon sublayers are already present on the fibers extracted from the preform prior to the SiC CVI­densification /15/.

A similar carbon sublayer in contact with Nicalon fibers has been already reported in glass-ceramic matrix CMC processed by hot pressing at a somewhat higher temperature (i.e. about 1200'C) /7-9/. According to R.F. Cooper and K. Chyung /8/ or P.M. Benson et a1. /16/, the formation of the C/Si02 couple, at the Nicalon fibre surface in such composites, can be predicted thermodynamically as the result of a chemical reaction occurring between oxygen from the matrix and SiC from the fiber and leading to the simultaneous formation of silica and carbon.

As shown in fig. 3, crack propagation (mode II) occurs, at/or near the BN/Si02 interface, a feature which could be due to : (i) a poor adhesion between pyrolytic boron nitride and amorphous silica or/and (ii) the weak bonding /17/ between the layers of the BN structure (which lie almost parallel to the interface in the vicinity of the BN­fibre boundary). As a matter of fact, the AES analyses performed on the failure surfaces of SiC {Nicalon)/BN/SiC-CVI composites have shown that the BN interphase is always found on the concave (matrix-side) parts of the failure surface (fig. 4 and 5).

4- CONCLUSION

TEM and AES analyses have shown that in SiC/BN/SiC composites the interphase deposited from BF3-NH3 precursor is made of turbostratic pyrolytic boron nitride with a N:B atomic ratio close to one and containing less than 5 at. % oxygen. Moreover, two thin layers of amorphous carbon and silica were found at the BN-Nicalon fiber boundary. These materials, whose occurrence has been already observed and thermodynamically justified for other CMC, are thought to be formed during the synthesis of the composite.

132

SiC-Matrix BN / Si02/ C/ Nicalon Fiber Fig. 1 The SiC-matrix/ pyrolytic

silica/ amorphous carbon/Nicalon fiber sequence as observed in TEM

BN interphase/ glassy in SiC/SiC composites,

random layers

aligned layers

BN Si02 C Fig 2 The microtexture of the BN-interphase in the vicinity of

the BN/silica interface, as observed from a bright field observation in TEM

N/B at. ratio oxygen d002 (EELS) (at. %) (A)

BN fiber coating 0.7 - 0.8 10 - 15 3.65 (before SiC-CVI)

BN interphase in 0.94 - 1.05 <5 3.51 SiC/BN/SiC

Table 1 : Chemical and structural features of BN used as an interphase material in SiC/SiC composites.

133

Fig. 3 : Crack propagation path at (or near ) the BN/silica interface in a SiC/SiC composite, as observed in TEM

Fig. 4 Failure surface of a SiC/BN/SiC composite

134

100

90

80

70

60

50 c. 'afI.

0 40 c(

30

20

10 o

5 10 15 20 25 30 35 40 SPUTIER TIME (MIN.)

Fig. 5 : AES depth profiling near the fiber-matrix boundary in a SiC/BN/SiC composite

REFERENCES

1- R.W. Rice, US Patent 4, 642, 271, Feb. 10, 1987 2- D.P. Stinton, A.I. Caputo and R.A. Lowden, Amer. Ceram. Soc.

Bull., 65/2 (1986) 347 3- B. Bender, D. Shadwell, C. Bulik, L. Incovarti and D. Lewis,

Amer. Ceram. Soc. Bull., 65/2 (1986) 363 4- R.N. Singh and M.K. Brun, Advanced Ceram., 3/3 (1988) 235 5- L. Grateau, N. Lob and M. Parlier, to be published in Proc.

14th Int.Conf. Science of Ceramics, Canterbury, UK, 1987 6- M.H. Rawlings, J.A. Nolan, D.P. Stinton and R.A. Lowden, Proc.

MRS-Symposium "Advanced Structural Ceramics", Vol. 78, (1986) 223, 7- R. Chaim and A.H. Heuer, Adv. Ceram. Mater., 2/2 (1987) 154 8- R.F. Cooper and K. Chyung, J. Mater. Sci., 22 (1987) 3148 9- J.J. Brennan,. in "Tailoring Multiphase and Composite Ceramics"

(R.E Tressler et al. eds.), Mat. Sci. Res. 20 (1986) 549 10- H. Hannache, R. Naslain and C. Bernard, J. Less-common Metals,

95 (1983) 221 11- H. Hannache, J . M. Quenisset, R. Naslain and L. Heraud, J.

Mater. Sci., 19 (1984) 202 12- F. Chris tin , R. Naslain and C. Bernard, Proc. 7th Int. Conf.

CVD (T. O. Sedwick and H. Lydin, eds. ), Los Angeles, p. 499, The Electrochem. Soc., Princeton, 1979

13- R. Naslain, J.Y. Rossignol, P. Hagenmuller, F. Chris tin , L. Heraud and J.J. Choury, Rev. Chimie Minerale, 18 (1981) 544

135

14- R. Naslain and F. Langlais, in "Tailoring Multiphase and Composite Ceramics" (R.E Tressler et al., eds) , Mat. Sci. Res. 20 (1986) 145

15- O. Dugne, S. Prouhet, A. Guette, R. Naslain, R. Fourmeaux, K. Hssein, J. Sevely, C. Guimon, D. Gonbeau and G. Pfister-Guillouzo, to be published in the Proc. of Euro CVD-7, Perpignan, France, 1989

16- P.M. Benson, K.E. Spear and C.G. Pantano,to be published in Proc. 12th Am. Conf. on Composites and Advanced Ceramics, Cocoa Beach, 1988

17- W. Sinclair and H. Simons, J. Mater. Sci. Let., 6 (1987) 627

MATRICES METALLIQUES METAL MATRIX

Chairmen: Pr M. G. BADER University of Surrey Dr K. SCHULTE D.F.v.L.R.

ABSTRACf

CAST FIBRE REINFORCED ALUMINIUM ALLOY MICROSTRUCTURES

R. TRUMPER, V. SCOTT"

Admiraffy Research Establishment Holton Heath - BH16 1JU POOLE DORSET - England

'University o( Bath Claverton Down, BA2 7 A Y BA TH - England

Examination of fibre-reinforced metal microstructures produced via a pressure infiltration route shows that for hypoeutectic alloys, primary metal dendrites nucleate and grow from within the interfibre regions, resulting in segregation of second phases to the fibre-matrix interface. If these intennetallics cannot be removed during a subsequent solution treatment then they may act as Griffith type flaws and substantially reduce the tensile strength of the composite.

INTRODUCTION

The successful use of a molten alloy route to fabricate a fibre-reinforced metal (FRM) must overcome a number of fundamental problems such as satisfactory infiltration of the fibre prefonn, wettability of the fibre surface and reactivity between fibre and matrix. Early attempts to fabricate FRMs via a liquid metal process attempted to overcome the poor wettability of common reinforcements (boron, silicon carbide, alumina and carbon) in aluminium, by either alloying the metal with elements such as lithium, magnesium, indium, lead and thallium /1,2/, or by the use of elevated melt temperatures to reduce the contact angle /3/. These approaches, however, usually resulted in rapid chemical attack of the fibre by the liquid alloy, and serious degradation of fibre properties.

Recently, much attention has been directed towards pressure assisted infiltration, where the molten matrix is forced into the fibre network using sufficient positive pressure to overcome the surface tension forces. This approach has a number of attractions. Firstly, no fibre/matrix reaction is necessary to produce a fully infiltrated composite and hence the maximum mechanical properties conferred by the fibres could be realised by the composite (assuming adequate shear transfer across the fibre/matrix interface). Secondly, fibre-reinforced components can be manufactured to near net-shape, so avoiding the very high secondary fabrication costs associated with traditional diffusion bonding techniques.

140

Finally, a variety of different fibres can easily be incorporated in the same component to allow the use of selected, or hybrid reinforcement.

This paper examines a number of different FRM systems manufactured using a liquid metal infiltration process developed at ARE 141 and considers some consequences of using the process to fabricate metal matrix composites.

RESULTS AND DISCUSSION

Figure 1 is an optical micrograph of a silicon carbide coated boron fibre (Borsic) incorporated in a 7018 matrix (Al-4.Swt% Zn-2.5wt% Mg alloy); the micrograph was taken using lead oxide film, multipath interference contrast 151. The presence of a second phase which has segregated to the fibre/matrix interface can be clearly seen and transmission electron microscopy, in conjunction with x-ray analysis and selected area diffraction techniques 16/, identified this phase as Mg2Si; no reaction between the fibre and matrix was found Hence, it may be deduced that such phases are formed by reaction of the magnesium with silicon impurities in the alloy, rather than with the silicon carbide coating of the Borsic fibre. The location of this second phase is powerful evidence for the way in which solidification-Of the matrix has occurred. It indicates that primary aluminium dendrites have nucleated and grown from the interfibre regions and converged on the fibre network which, in turn, has acted as a physical barrier to pin dendrites and so limit the size of the dendrite arm spacing.

We have investigated a number of other FRM systems in order to study how the mode of solidification described above may be influenced by different alloy types and the physical or chemical nature of the reinforcement. Figure 2 shows a Nicalon fibre reinforced Al-lwt% Fe alloy. Segregation of an iron-rich phase (FeA16) around most of the fibres is Clearly apparent, and the phase spans some of the matrix rich areas. This microstructure is also consistent with the nucleation and growth of aluminium dendrites in the interfibre regions and the consequent rejection of solute elements. As solidification proceeds, the remaining liquid is progressively enriched in solute until the eutectic composition is reached and then second phases are formed in the last areas to solidify, those commonly adjacent to the fibre-matrix interface. The microstructural scale of this system is similar to but an order of magnitude smaller than that of the Borsic AI-Zn-Mg composite discussed above. These results clearly demonstrate the physical influence that the reinforcement has on the scale of the solidified structure under essentially similar thermal conditions.

The way in which molten metal interacts with a fibre bundle has been studied by a number of workers. Fukunaga & Ohde m examined the passage of the melt front through a network of cold fibres (ie fibres at a temperature below that of the alloy solidus) and developed a model whereby heat was transfered from the melt to the fibres. One result was the build up of a solid metal layer on each fibre which effectively choked off the capillary channels so preventing any further progress of the melt front. This effect would lead to incomplete infiltration and would obviously be undesirable.

The rate at which the capillary channel is reduced in width during the infiltration process may be assessed as follows: consider frrst the local heat exchange time, TI' between fibre and melt

141

( 1 )

where D is the fibre diameter and ex is the fibre thermal diffusivity. For most FRM

systems of interest this gives an exchange time of'" 10-3 sec, which implies that heat is transfered rapidly from liquid metal to fibre when they come into contact. If a square­packing geometry is assumed, see figure 3, there are two possibilities:

I) If the resultant composite temperature, Tc' is greater than T m' the matrix liquidus, then;

2) IfTc = Tm then a proportion of solid matrix, V s' has been formed;

where V f and V s are volume fractions of fibre and solid matrix, Cm and Cf are specific heat capacities of matrix and fibre, T a' T f and T m are the temperatures of alloy, fibre and

alloy liquidus, Lm is the alloy latent heat of fusion and Pm' Pf are the densities of matrix and fibre.

Formation of a proportion V s of solid matrix leads to a reduction in the capillary channel width P according to:

(4)

Thus when P = 0, infiltration of the preform would cease. However, the microstructures of fully infiltrated composites show little or no evidence for the formation of solid layers of matrix around individual fibres. This implies that reliance on the contribution of alloy latent heat of fusion for local heating of the fibre network is unlikely to result in a successful composite, especially when the random packing of fibres in a real system is taken into account, see figure 2.

Figure 4 shows the relationship between V s and melt superheat at three different fibre preheat temperatures for a 50% V f Nicalon fibre! AI composite. These data indicate that to avoid problems of incomplete infiltration as well as minimising infiltration pressures, Tc should be slightly greater than the alloy liquidus.

As has been pointed out, the use of a pressure infiltration route with an alloy matrix leads to segregation of second phases around the fibres, not all of which can be removed by subsequent solution treatment. This gives rise to the possibility of brittle intermetallics acting as flaws in an analogous manner to the fibre-matrix reaction products described by

Ochiai et al/8/. These workers assumed that fibre strength Of decreased with increasing

reaction zone width (c) according to a Griffith relationship:

(5)

142

where Er is the fibre elastic modulus and Gc * is the critical strain energy release rate for

the fibre. Tensile strengths have been calculated for a 25% volume fraction Safimax/AI composite using Zweben's statistical failure criterion /9/ to predict the lower bound at

small values of c, taking Gc * = 28 Jm-2, see figure 5. The data indicate that this system

should be insensitive to intermetallics of less than = 0.7 Ilm thick. If this prediction is compared with the experimentally measured tensile strength of such a system where FeAl6 has segregated to the fibre surface, a reasonable correlation can be found.

It is tempting to suggest that if fibre and melt temperatures could be controlled with sufficient accuracy, it might be possible to form a very thin layer of relatively pure matrix as a skin around each fibre and so avoid the problem of intermetallics segregating there and acting as flaws.

Summarising, we have shown that a number of distinctly different FRM systems, based on hypoeutectic matrices, exhibit very similar microstructural characteristics which can be accounted for by the nucleation and growth of primary aluminium dendrites within the interfibre regions. For relatively rapid cooling rates, the actual scale of the microstructure is strongly influenced by the physical presence and spatial arrangement of the reinforcement. As in conventional hypoeutectic casting metals, alloying elements and impurities tend to segregate to the interdendritic regions during the final stages of solidification which tends to occur at the reinforcement. Thich implies that even a relatively small amount of impurity may strongly influence the mechanical properties of the composite.

ACKNOWLEOOEMENTS

To SERC for Financial support. © Controller, HMSO, London, 1988.

REFERENCES

1. Delanney F., Froyen L. and Deruyttere A., 1. Mat. Sci., 22, 1987, 1-16 2. Kimura Y., Mishima Y., Umekawa S. and Suzuki T. 1. Mat. Sci., 19, 1984,

3107. 3. Warren R and Andersson C.H., Composites, 15,2, 1984, 101-111. 4. Trumper RL., Sherwood PJ. and Clifford A.W. Proc. Conf.,"Materials in

Aerospace", Royal Aeronautical Society, London 2-4 April, 1986, Vol II, 249-289. 5. Bennett E.G, 1. Met. & Mat., 3,1987, 278-279. 6. Saggese M.E., Scott V.D. and Trumper R.L., Mat. Sci & Tech., 4, 1988,871-

875. 7. Fukunaga H. and Ohde T., "Progress in Science and Engineering of Composites. ",

Eds. T. Hayashi, K. Kawata and S. Umekawa, ICCM-IV, Tokyo, 1982, Japan Society for Composite Materials, 1443-1450.

8. Ochiai S., Osumura K. and Murakami Y., "Progress in Science and Engineering of Composites.", Eds. T. Hayashi, K. Kawata and S. Umekawa, ICeM-IV, Tokyo, 1982, Japan Society for Composite Materials,1331-1338.

143

9. Zweben c., Smith W.S. and Wardle M.W., "Composite Materials: Testing and Design.", ASTM STP 674, American Society for Testing and Materials, 1979,228-262.

1---1 25 J.I.m

Fig. 1 Al-Zn-Mg alloy reinforced with.Borsic fibre, optical micrograph, lead oxide fIlm interference contrast.

1--1 25J.l.ffi

Fig. 2 AI-I wt. % Fe alloy reinforced with Nicalon fibre, optical micrograph, Kellers etch.

fibre

solid matrix

Fig. 3 Square fibre packing geometry showing layer of solid matrix on fibres.

144

0.4

'9 ~

0.3

;g ~ 0.2

'Ci

~ "0 > 0.1

0.0

650

Fibre Temp. °C

D 400 • 500 .. 600

700 750 800

Melt temperature °C

Fig. 4 Calculated volume fraction of solid matrix for a 50% V f Nicalon fibre reinforced aluminium composite as a function of fibre and melt temperatures.

1000

0:1

~ 800 . -5 b.O c:: 600 ~

. \ Statistical failure criterion '" £

400 ..... '"

! 200

- ~ ~ Experimentally detennined strength

~ Griffith criterion

0 . o 2 4 6 8 10 12

Intennetallic width Jlm

Fig. 5 Predicted composite failure stress for a 25% V f Safimax fibre reinforced aluminium alloy as a function of intennetallic width.

THE INFLUENCE OF THERMAL CYCLING ON THE PROPERTIES OF SI3N4 WmSKER REINFORCED

ABSTRACT

ALUMINIUM ALLOY COMPOSITES

Y. NISHIDA, M.H. MASARU, M.Y. NAKANISHI

Government Industrial Research Institute Nagoya Hirate-cho, Kita-ku, NAGOYA 462 - Japan

Thermal cycling tests were performed on the Si3N4 whisker/99.9% aluminum and /AC8A alloy composites obtained by squeeze casting, foll­owed by tensile tests, observation of fracture surfaces and microstru­ctures to examine the fatigue of matrices. Little change in tensile strength was observed for the specimen subjected to 1,000 cycles, ex­cept T6 treated AC8A alloy matrix composite. After 500 cycles, the effect of T6 treatment was lost. However, fracture surfaces showed somewhat brittle surface, which should be caused by locall plastic flow of matrice.

INTRODUCTION

Ceramic whisker reinforced aluminum alloys are prom~s~ng materials for engine parts, because their properties, especially at elevated tem­peratures are excellent/1,2,3/. In addition, three-dimensional streng­thening and plastic deformation are possible for those composites.

Those composites will be heated and cooled cyclically when those composites are put to practical use. When those composites are exposed to thermal cycling, with which the stress exceeds the elastic limit of the matrix, the matrix metal will be deformed locally (plastic flow) and degraded. Then, voids might be formed at the interface between matrix metal and whiskers, due to the mismatch of their thermal expan­sion coefficients. Although the research on the influence of thermal cycling on whiskers reinforced aluminum alloys are of importance from the practical viewpoint, there are few papers/4/.

On the other hand, the fabrication of Si3N4w reinforced aluminum alloys was performed and good mechanical properties which are compara­ble to SiCw reinforced aluminum alloys were revealed/5/. Then, this

146

study was carried out to make clear the effect of thermal cycling on the tensile strength, microstructure and fracture surface of Si3N~w reinforced aluminum alloy.

I - EXPERIMENTAL PROCEDURE

Si3N~w reinforced 99.9% aluminum and AC8A alloy (JIS alloy; Al-12% Si-1% Cu-1% Mg) composites were fabricated by squeeze casting. Si3N~ whisker used for the present study was a-type single crystal and had 0.1 to 1.6 um in diameter with aspect ratio of 20 to 200. Preforms made by an aspiration method were used for the fabrication of the com­posites. Preforms and the mold were preheated up to about 1023 and 573 K, respectively. The pressure applied during casting was 100 MPa.

Tensile specimens were made from the composites, Vf (volume frac­tion) of which was about 15 %. The diameter of the tensile test speci­mens at gauge marks was 6 mm. Some of tensile test specimens of AC8A alloy matrix composite were made after T6 heat treatment.

The schematic representation of the apparatus for thermal cycling is shown in Fig.1. An infrared-ray furnace was used for heating. A tensile test specimen was placed in a quartz tube in the center along the axis of the cylindrical furnace, and held by pushing with three bolts at both ends of specimen to reduce heat transfer toward the stai­nless tube. The specimen were cooled at both sides of specimen by blo­wing compressed air and the temperature control was performed by PID controller, detecting the temperature of specimen with a thermo-couple, which was welded directly to the specimen. A thermal cycle consists of four parts: linear heating, holding at maximum temperature, cooling only by blowing air and holding at minimum temperature. A typical ex­ample of thermal cycling is shown in Fig.2. The maximum and minimum temperatures were 673 K (or 748 K for several cases) and 373 K, respec­tively. One cycle time was 10 minutes, and tests were performed up to 1,000 cycles.

Microstructures before and after thermal cycling were observed by optical microscopy, EPMA. Fracture surfaces after tensile tests were also observed by SEM.

II - EXPERIMENTAL RESULTS AND DISCUSSION

2.1. Mechanical properties

Tensile strength after thermal cycling is shown in Figs.3 and 4. The tensile strength of Si3N~w/99.9% Al before thermal cycling was 240 MPa. Little change in tensile strength within 1,000 cycles was obser­ved. Tensile strength for as-cast specimen of Si3N~w/AC8A alloy com­posite before thermal cycling was 320 MPa and did not change within 1,000 cycles except the T6 treated sample. The tensile strength of T6 treated AC8A matrix composite decreased with increasing number of cycle and the effect of the heat treatment was almost lost within 500 cycles. The maximum temperature of 748 K was applied to some specimens to exa­mine the effect of the maximum temperature on the tensile strength. As can be seen from Figs.3 and 4, the effect of maximum temperature on the strength was little.

147

2.2. Fracture surface

Fracture surfaces of specimens after tensile test observed by SEM are shown in Fig.S for 99.9% Al matrix and in Fig.6 for AC8A matrix composites. Many dimples can be seen in the surface of the uncycled 99.9% Al matrix composite. After thermal cycling of 1,000 times, the fracture surfaces show that the composite became somewhat brittle and the size of dimples are ununiform. In the case of AC8A matrix compo­site, it is difficult to recognize the difference of fracture surfaces between before and after thermal cycling, though the composite should have become somewhat brittle. The similar change in fracture surface was observed in the case of SiCw reinforced aluminum alloys subjected to 1,000 cycles/6/.

2.3. Microstructure

Microstructure of AC8A matrix composite (T6 treated) observed by the optical micrscopy is shown in Fig.7, comparing with T6 treated AC8A alloy. Silicon crystals in the composites are finer than those in T6 treated AC8A. It is difficult to observe nickel base intermeta­llic compound in the composite before thermal cycling. However, many intermetallic compounds (white crystals) appeared after 500 cycles. Those crystals seems to affect the tensile strength of the T6 treated composite.

The difference in dimension of specimens between before and after 1,000 cycles was also measured, and it was less than errors in measur­ments, though Patterson et al./4/ obtained the change in the case of SiCw/2021 alloy composite fabricated by powder metallurgical method. This result suggests that dimensional change, thermal ratchet, depends on the fabrication methods of the composites and on the distribution of whiskers. However, local 1 plastic flow of matrix should occur, due to the mismatch of thermal expansion coefficients, when the amplitude of thermal cycling is so large that the internal stress exceeds the elastic limit of the matrix. The locall plastic flow causes the fatigue of matrix.

III - CONCLUSION

The effect of thermal cycling on the properties of Si 3 N4 w/99.9%Al and Si3 N4 w/AC8A alloy fabricated by squeeze casting was examined. The tensile strength did not change within 1,000 cycles. However, the effect of thermal cycling on the matrix appeared in the fracture sur­faces of tensile test specimen, although dimensional change of the specimen subjected to 1,000 cycles was negligible. The indication of fatigue of matrix seems to be due to locall plastic flow of matrix.

REFERENCES

1 - D. Webster, Met. Trans. A13A(1982) 1511. 2 - T. Imai, Y. Nishida, M. Yamada, H. Matsubara and I. Shirayanagi,

J. Mater. Sci. Lett., 6(1987) 343. 3 - T. Imai, Y. Nishida, M. Yamada, I. Shirayanagi and H. Matsubara,

ibid., 6(1987) 1257.

148

4 - W. G. Patterson, M. Taya, Proceedings of 5th ICCM, (1985) 53. 5 - H. Matsubara, Y. Nishida, M. Yamada, I. Shirayanagi and T. Imai,

J. Mater. Sci. Lett., 6(1987) 1313. 6 - M. Nakanishi, Y. Nishida, H. Matsubara, M. Yamada, Y. Tozawa and

M. kato, Reprints of 75th Conference of Japan Institute of Light Metals, (1988) 37.

Quartz l ube

Compressed air

Stainless tube Stainless ring

Thermo-couple

Fig. 1 Schematic representation of apparatus for thermal cycling.

;::., 700 400 ,..----r---r-----,.-----,..---r-.....,

<II

; 600 OJ

~ 500 <II 0-

@ 400 E-

10 20

Time (min)

Fig . 2 Example of ther mal cycling .

<0

~ 300

.s::::. 0-. +.J g'200 Q) ~ +.J

Vl100 Q)

Vl

0_ -0

maximum ma t erial temp .(K)

673 748 0 • Si 3 N. /AI

~ a L-____ ~ ____ ~ ____ ~ ____ ~ ____ ~~

I- 0 200 400 600 800 1000 Number of cycles

Fig. 3 Relationship between tensile strength of Si3N4 whisker/99.9% Al alloy and number of thermal cycling.

600 maximum temp. 673 748(K) .

material

100 ~

o L..._..L..._-,-__ ~I_~ I_~I--

o 200 400 600 800 1000 Number of cycles

Fig. 4 Relationship between tensile strength of Si 3 N4whisker/ AC8A alloy and number of thermal cycles.

(a) I 5/Jm I

149

Fig. 5 Fracture surfaces of Si3N4 whisker/99.9% Al composite observed by SEM, (a) as-cast, (b) subjected to 1,000 cycles.

150

(a) Fig. 6 Fracture surfaces of Si3N4 whisker/AC8A alloy composite observed by SEM, (a) as-cast, (b) subjected to 1,000 cycles.

(a) ( b) (c)

Fig. 7 Microstructure of AC8A matrix composite and AC8A alloy observed by the optical microscopy, (a) as-cast composite, (b) composite subjected to 500 cycles, (c) as-cast AC8A alloy.

FRACTURE OF AL·SICW MET AL MATRIX COMPOSITES

Abstract:

D. CHAMBOLLE, D. BAPTISTE, P. BOMPARD

Ecole Centrale de Paris· Laooratoire des Materiaux Grande voie des Vignes . 92290 CHATENA Y MALABRY . France

Reinforcement of silicon carbide whiskers improved significantly the elastic

and strength behaviours of aluminium alloys. Unfortunately, the increase of the

volume fraction of whiskers is responsible for a very low ductility of these

M.M.C. and the sensitivity to stress concentration is very high. In order to use

these materials in some structural components, we need to obtain a fracture

criterion for the 2124 Al + 20% SiCw at room and high temperatures (i.e.20°C to

350°C). After a tension and fatigue crack growth characterization, a damage

micromechanics modelisation based on the microstructural characteristics such as

the aspect ratio and the volume fraction distributions is established.

Tension and fatigue testing:

Tensile and fatigue tests were performed on a 2124 Al T6 + 25% SiCw

extruded composite provided by Arco, extruded by Pechiney. Tensile specimens

were machined along and across the fiber axis of the composite. They were

submitted to high rate heating and tested in traction. This material exhibits higher

tensile modulus than Titanium alloys up to 350°C and a very high strength up to

200°C. But, there is no more reinforcement effect at high temperature related to a

2024 T6 alloy (this will be discussed later). Fatigue crack growth tests were

performed at room temperature with compact tensile specimens (C.T. 12 mm)

along and perpendicularly to the fibers axis. We were unable to propagate a crack

152

in a plane across the fibers axis (Le. with a C.T. specimen machined along the

fibers axis, L-T direction). The behaviour of this M.M.C. is not too different from

the matrix one for the lowest values of ~KeI For higher values of ~K~ crack

growth rates are much higher, due to the very low value of KIc (6 MPa ..Jm). The

fracture toughness was measured with the C.T. specimens for the T-L mode and,

with three points bending specimens (stress corrosion precracked) for the L-T

mode (Kl:= 20 MPa ..Jm).

Mechanical behaviour of the interfaces;

The most discussed fact, concerning the silicon carbide whiskers reinforced

Aluminium, is the good behaviour of the fiber/matrix interfaces. Some authors /11

are convinced that there is a very strong chemical bond (due to a reactivity of Al

and SiC, for instance), some others /2/ have shown that there was not any

chemical reaction between the Al matrix and the SiC up to the melting point of the

matrix, except for very long times of interaction /3/.

Assuming that great reinforcement effect is related to a good stress transfer

between the fibers and the matrix, the drastic fall of the strength around 200°C led

us to consider a decrease of this stress transfer at this temperature. Moreover, at

this temperature, we can observe pulled out fibers on the fracture surfaces.

Different computations (Fig 1) and (Fig 2) and X ray measurement /4/ of the

thermal stresses showed that their level decreases to zero at about 250°C. These

residual stresses arise from the mismatch of strains between the fibers and the

matrix during the cooling which follows the heat treatment of the material. This

mismatch is mainly due to the difference of the coefficients of thermal expansion

between the matrix (23 1O~-1 at room temperature) and the fibers (5 1O~-~. If

we consider that the good stress transfer is also improved by the very rough

surface of the fibers, we can explain the fall of the strength from 200°C with the

rapid decrease of the radial residual stresses around the fibers.

Damage micromechanics modelisation :

In order to establish a damage criterion at the scale of a fiber, we first observe

the damage process occuring around the fibers using in-situ tension tests ( tension

test in a scanning electron microscope). To modelize the damage process, we

calculate the local stress-strain field in and around a fiber using a finite element

153

In-situ tension tests. Quantitative m.etaUognmhy :

Tension tests were conducted in a scanning electron microscope, at room and

high temperature. The pictures were recorded and post-processed with a picture

analyser. Fig 3 shows a 30mm crack slowly propagating across the specimen. We

can identify a two steps mecanism of failure :

During the first step, the fiber/matrix interface yields at the fiber tip.

The second step consists in the failure of the matrix. It fails from fiber tip to

fiber tip. The cracks often follow the fiber/matrix interfaces. The general

orientation of the crack is normal to the highest principal macro stress (1:). We

have determined, along the crack the local values of some microstructural

parameters (Le. fibers aspect ratio, volume fraction of whiskers and raw particles,

misorientations) (Fig 4).

These local values were compared to the global ones in order to identify the

most effectives parameters related to the fracture of the material.

Damage micromechanics modelisatiQn :

To correlate the damage process observed at the fiber scale, finite element

calculation were performed on a basic cell composed of the fiber (aspect ratio Vd),

the matrix in proportion of the local volume fraction Vf, the overall surrounded by

the equivalent homogeneous material related to these previous features. Residual

stresses, induced by the cooling during the elaboration process, were first

calculated. (Fig 2). They are a function of the temperature amplitude between the

heat treatment temperature and the service one. Then, applying a macro-stress

field :E on the basic cell, we determine the local stress-strain field in the fiber and

the matrix. Knowing a local damage criterion in the matrix (cavity growth or

plastic instability), we determine the macroscopic stress field:ED for a given local

damage state. By changing the fiber aspect ratio, Vd, or the local volume fraction

of fibers, Vf, in the basic cell, we obtain the scatter of the macroscopic strength

tensor as a function of these microstructural parameters (Vd, Vf). This scatter can

be compared with the one obtained on strength of notched specimen in tension /5/.

This statistical distribution of stresses to failure can be used in a finite ele!Dent

analysis of a structure to determine its probability of failure. The algorithmic

procedure can be split up as follows :

154

1°) The structure is meshed; the finest elements of the mesh (located in the critical

zones) are sized in accordance with the volume of the basic cell used for the

determination of the criterion,

2°) An elasto-plastic computation of the stress field is performed. In this step, we

have to take into account the transverse isotropy of the material,

3°) The probability of failure, pi, of each element is determined with the

distribution of stresses to failure ~ as determined previously,

4°) The probability of failure of the structure is given by :

Pi is the probability of failure of the ilh element, Vi is its volume; Vo is the

volume of the elementary cell.

Conclusion:

This micromechanics approach, based both on finite elements calculation and

on experimental observation allowed us to describe the damage state of a

2124 Al + 20% SiCw from the characteristics of the components. The

modelisation showed the influence on the strength of the different fiber aspect

ratios and local volume fraction Vf and the residual stresses induced by the

elaboration process. It leads to a scattered damage criterion which can be used in

structural components calculation in order to predict the failure.

This work has been supported by the Societe Nationale Aerospatiale.

1. Flom, Y., Arsenault, R. J., "Fracture of SiC/AI composites", (1987), Proc.

ICCM VI & ECCM 2, London.

2. Nutt, S. R., "Interfaces and failure mechanisms in AI-SiC composites" in

Interfaces in metal matrix composites, (Dhingra and Fishman ed), (1986).

3. lC. Le Flour, "Elaboration d'alliages d'Aluminium et de composites base

Aluminium par metallurgie des poudres", PhD Thesis E.N.S.M.P., (1987).

155

4 PJ. Whiters, D. Juul Jensen, H. Lilholt and W.M. Stobbs,"The evaluation of

internal stresses in a short fibre metal matrix composite by neutron diffraction",

(1987), Proc. ICCM VI & ECCM 2, London.

5 R.A. Hunt and L.N. Mc Cartney, "A new approach to Weibull's statistical

theory of brittle fracture", (1979), Int. J. of Fracture, Vol. 15, No 4

Fig 1 Evolution of the

residual stress a in the

aluminium, computed with a

two phases mono-axial model,

during the thermal cycle :

450°C -> 20°C -> 450°C.

Note the very low value of ali

around 250°C during the

heating.

Equ/v(Jlent homogeneous

aMPa

300

100

-100

-200

-300

Fig 2 : Three phases

model used for the

computation of the

localization of

stresses and strains

around the fiber.

Residual radial

stresses arr for the

thermal cycle

450°C -> 20°C

m(Jter/(J/. H+t++~f+++-+--f

H(Jtr/x.

FIber.

500

a, MPa

o.

~~~N-50. -100.

-150. -100 .

156

Fig 3 : Crack path across a composite submitted to mono axial loading along the fiber

axis in a scanning electron microscope.

100% Pc •• • • • .. •• • • 2

•• So =16 1lm

~

•• • • !l Vf ~ 0%

- 1 00,00% O,OO:C 100,00% 200,00:1:

Fig 4 : Probability of having a volume element with a greater or lower volume fraction

of silicon carbide than the mean value.

PHYSICOCHEMICAL ASPECTS OF THE K2ZrF6 PROCESS ALLOWING THE SPONTANEOUS

INFILTRATION OF SIC (OR C) PREFORMS BY LIQUID ALUMINIUM

ABSTRACT

S. SCHAMM, J.P. ROCHER', R. NASLAIN"

Laboratoire de Chimie du So/ide du CNRS 351 cours de /a Liberation - 33405 TALENCE - France

'SEP - BP 37 - 33165 ST MEDARD EN JALLES - France "Laboratoire des Composites Thermostructuraux

Europarc, 3 avenue Leonard de Vinci - 33600 PESSAC - France

K2ZrF6 enhances the wetting ability of SiC (or C) fibers by liquid aluminum alloys at low temperatures due to (i) a dissolution of the alumina film by fluoride species, (ii) a local evolution of heat related to exothermic reactions and (iii) the formation of a new liquid phase. The K2ZrF6 process resul ts in a lowering of the UTS reinforcements due to an activation of the kinetics of formation of A14C"1 crystals at the fiber surface. Based on a theoretical modelization, it is established that the impregnation ability of SiC­SiC preforms by aluminum is significantly improved by a K2ZrF6 treatment applied to the preforms prior to casting.

INTRODUCTION

Among the processing techniques which have been reported for aluminum matrix composites (AMC) , those based on casting (e.g. gravity casting or squeeze casting) seem to be particularly suitable for volume production of near net shape parts at an acceptable cost. However, the impregnation of a fiber preform by liquid aluminum alloys is often difficult due to : (i) the poor wetting of most ceramic fibers (e.g. carbon or SiC) by the liquid alloys and (ii) extensive chemical reactions taking place during processing between the fibers and the matrix. Casting under high pressure (e.g. squeeze casting) applied to alloys of optimized compositions is presently the most efficient way to overcome these difficulties. However, it requires pressing equipments of large capacity when applied to parts of large sizes.

J.P. Rocher and al. have recently shown that the wetting ability of a ceramic fiber preform by liquid aluminum alloys could be significantly improved when the preform is pre-treated with an aqueous solution of K2ZrF6 according to a very simple procedure, prior to casting /1/. In its principle, the process is based on an impregnation

158

of the preform by an aqueous solution of K2ZrF6. During the following drying step, microcrystals of K2ZrF6 are formed at the fiber surface which were found to promote the wetting of both carbon or SiC-based fibers by aluminum alloys at low temperatures. As a result, casting can be performed according to the classical gravity casting procedure (Le. the use of high pressures is not necessary). The efficiency of the process was demonstrated on different kinds of fiber preforms including porous 2D-SiC (Nicalon)/SiC (CVI) preforms /1, 2/.

The aim of the present contribution is to present and discuss the scientific bases of the now so-called K2ZrF6 process and to study its effect on the strength of SiC-based reinforcements.

1- WETTING ENHANCEMENT DUE TO K2ZrF6

The spontaneous wetting (9 < 90·) of carbon or SiC-based ceramics, pretreated according to the'K2ZrF6 process, by aluminum has been established both from impregnation experiments run on SiC (or C) fiber preforms and from sessile drop experiments performed on plane SiC substrates.

For untreated carbon or SiC substrates, the contact angle usually reported in the literature for temperatures ranging from 700 to 800·C, is of the order of 160·, for pure aluminum /3/. It falls to 60-75· when the substrates have been pre-treated according to the K2ZrF6 process, as shown in fig. 1 /4/. Furthermore, it was observed that 9 decreased as the amount of K2ZrF6 deposited on the substrates was increased. It is noteworthy that these low 9 values correspond almost to those recently reported by V. Laurent et al. for SiC single crystals and aluminum free of any alumina contamination /5/.

The mechanisms that could explain the wetting enhancement due to K2ZrF6 might have different origins, namely : (i) a dissolution of the alumina film which is usually present at the surface of liquid aluminum, (ii) a local temperature increase due to exothermic reactions, and (iii) the formation of a new liquid phase.

1.1- Dissolution of the alumina film

When liquid aluminum flows within the pore network of a fiber preform, which has been pre-treated with K2ZrF6, the two following chemical reactions are thought to take place :

3K2ZrF6 + 4 Al .... "6KF, 4AIF3" + 3Zr 3Zr + 9Al .... 3A13Zr

[1] [2]

since both of them are characterized by a Gibbs free energy variation of _ 284.4 and _ 487.2 kJ.mol-1 respectively, at 7oo·C. The fluoride mixture "6KF, 4AIF3" (or K3AIF6 + 3KAIF4) which is formed according to [1] may dissolve tne thin layer of alumina which is usually present at the liquid aluminum surface and prevents the spontaneous wetting of the ceramic (it is well known that cryolite Na3AIF6 is used to dissolve alumina in the electrolytic processing of aluminum). A detailed thermodynamic study of the "6KF ,4AlF3" /a-A1203 system has indeed shown that under our experimental conditions (Le. an alumina film 10 nm thick and a K2ZrF6 deposit of a few mg/cm2), corresponding

159

to about 1 mol% alumina, alumina may be dissolved at a temperature as low as 520·C /6/.

1.2- Local temperature increase

The reactions [1] and [2] are strongly exothermic (i.e. at 700·C, H [1] = - 306.4 and H [2] = - 487.2 kJ. mol-1). As a result, a

local temperature increase may take place near the fiber-matrix interface. Based on a model which will be discussed below (see § 3 and /8/), it has been established that the temperature may reach about 800· C during the impregnation process in 2D-SiC (Nicalon) /SiC (CVI) preforms pre-treated with K2ZrF6 (deposit of a few mg/cm2) and preheated to 650·C when reactions [1] and [2] take place.

1.3- Formation of new liquid phase

From the calculated vertical section of the KF_AIF3-AI203-K20 given in fig. 2, it appears that the "6KF, 4AIF3" fluoride mixture gives rise to an homogeneous liquid phase at the SiC (or C)/AI interface, when it is mixed with alumina at a temperature of at least 800·c (i.e. the temperature thought to be reached at the interface due to the exothermicity of reactions [1] and [2]) /6/.

The dissolution of the alumina film cannot alone fully explain all the features of the sessile drop experiments shown in fig. 1. It is thought that the formation of a new liquid phase at the fiber­matrix interface (modyfying both 0LV and osLl associated with the local temperature increase due to the exothermicity of [1] and [2] may also contribute significantly to the low 9 values which have been measured for the K2ZrF6 pre-treated SiC substrates.

2- EFFECT OF THE K2ZrF6 PROCESS ON THE REINFORCEMENT STRENGTH

The effect of the different steps of the K2ZrF6 process on the residual tensile strength of SiC reinforcement has been studied on SiC-CVD filament (d = 100 ~m, from SIGMA). The tests were performed at room temperature on lots of 40 filaments. The residual UTS data were analyzed on the basis of a Weibull two-parameter distribution. The reaction products at the filament surface were analyzed by scanning electron microscopy (SEM) and X-ray electron probe microanalysis (EPMA) •

2.1- Residual UTS Weibull plots

The first step of the K2ZrF6 process (i.e. the deposition of K2ZrF6 microcrystals from a hot aqueous solution) does not modify the UTS of the filaments. After an annealing treatment performed under vacuum at 650·C for 15 min, the mean residual tITS reaches a constant value equal to about 80% of that of the as-received filaments when the fluoride amount becomes large enough to form a continuous layer at the filament surface (i.e. about 5 mg.cm-2) (fig. 3).

On the contrary, when the filaments have been embedded in aluminum (by hot pressing between two aluminum foils, at 700·C for 5 min under a uniaxial pressure of 2 MFa and vacuum), a significant drop in UTS was observed on the filaments extracted from the matrix by a

160

NaOH treatment. The mean residual UTS was found to decrease as the amount of K2ZrF6 deposited at the filament surface was increased. When this amount was limited to a few mg/cm2 (i.e. to values high enough to resul t in good impregnations of complex porous fiber preforms), the mean residual UTS was still higher than 50% of the initial filament UTS.

2.2- Analysis of the degradation mechanism

The phenomenon which may be responsible for the strength lowering of the SiC-filaments has been identified from experiments performed under severe interaction conditions. Liquid aluminum was poured on SiC filaments coated with a thick K2ZrF6 deposit (Le. 8 mg/cm2) and the whole was maintained for 15 min at BOO·C. After dissolution of the matrix in a methanol solution added with brome the filament appeared to be heavily coated with hexagonal crystals as shown in fig. 4. On the basis of EPMA data, these crystals were identified to Al4C3. It is noteworthy that an experiment performed, under the same conditions, on a SiC-filament which has not been pre-treated with K2ZrF6, resulted in almost no formation of Al4C3 (in agreement with the data given in fig. 3 for the treated and untreated filaments extracted from aluminum).

It thus appears that the pre-treatment of SiC-reinforcements with K2ZrF6 may enhance the kinetics of formation of Al4C3 (which is known to result from the chemical reaction between SiC ana aluminum). This feature can be explained on the basis of the following considerations: (i) the K2ZrF6 process, by enhancing wetting, permits a good contact between SiC and aluminum, which is not the case otherwise (fig. 1), (ii) it dissolves the alumina film which may act as a diffusion barrier in untreated fibers and finally (iii) it increases the formation rate of Al4C3 by raising the interface temperature due to the exothermicity of reactions [lJ and [2J.

Therefore, the strength reduction of the SiC reinforcements observed in the K2ZrF6 process does not seem to be related to a chemical attack by fluoride species but simply to an activation of the kinetics of growth of Al4C~ due to the cleaning of the interface and to a local temperature 1ncrease. Obviously, the use of a soft interphase (such as carbon in the SCS-2 AVCO SiC filament) will suppress the notch effect, due to the failure of the brittle Al4C3 layer, on the filament surface /7/.

3- MODELIZATION OF THE IMPREGNATION OF PRE-TREATED PREFORMS

3.1- The impregnation model

Recently, J .M. Quenisset et a1. have presented a model for the impregnation of a ceramic fiber porous preform by liquid aluminum alloys /4, 8/. Their model is based on the equations of mass and heat transfers applied to a semi-infinite 1D preform made of cylindrical fibers in hexagonal packing. The impregnation ability of the preform is characterized by (i) the maximum impregnated depth Zpmax and (ii) by the pressure which has to be applied to the liquid metal at the preform surface to result in an impregnation at a depth Zp. As a matter of fact, Zpmax is the value of Zp for which the temperature of the liquid is equal to its melting point.The model has been used here

161

to compare the impregnation abilities of SiC (Nicalon) based preforms supposed to have been pre-treated or not according to the K2ZrF6 process. It was assumed that the mean diameter of the SiC Nicalon tows within the preform was 360 pII, after the SiC-CVI consolidation step (overall residual porosity : 25%). The temperatures of the liquid aluminUII and of the fiber preform were assUlled to be 800 and 650 0 C, respectively. Reactions [1] and [2] were assumed to take place at the fiber/metal interface as the liquid metal front propagates within the preform.

3.2- Results

The results of the calculations are given in table 1. It clearly appears that the preforms behave differently with respect to impregnation mainly due to .(i) their different wetting abilities and (ii) the occurence of exothermic reactions (reactions [1] and [2]) within the preform pre-treated with K2ZrF6. The capillary pressure, which is proportional to _ 0 cosO (0 : interfacial tension of the liquid ; 0 : contact angle), is negative for the preform pretreated wi th K2ZrF6 (0 = 75 0

) and positive for the untreated preform (0 = 160 0

). As a result, it decreases the overall pressure, which is necessary to impregnate the preform at Zp, to values lower than the metallostatic pressure used in gravity casting. Furthermore, the heat evolution due to [1] and [2] prevents the solidification of aluminum to take place as the metal flows in the preform (such solidification being indeed observed in the untreated preform).

ACKNOWLEDGEMENTS

The authors are indebted to C. Bernard, N. Eustathopoulos and R. Fedou for their valuable assistance and advices in the thermodynamics, wetting and impregnation modelization studies.

REFERENCES

1- Rocher J. P ., Quenisset J. M., Pailler R. and Naslain R., European Pat., 80901 204-4

2- Rocher J.P., Macke T., Quenisset J.M., Naslain R. and Cotteret J., Proc. JNC-5 (C. Bathias and D. Menkes, eds.), pp. 343-360, Pluralis, PariS, 1986

3- KOhler W., Aluminium, 51 (1975) 443-444 4- Rocher J.P., Thesis * 888, Univ. Bordeaux, 1986 5- Laurent V., Chatain D. and Eustathopoulos N., J. Mater. SCi, 22

(1987) 244-250 6- Schamm S., Rabardel L., Grannec J., Naslain R. and Bernard C.,

submitted to Calphad 7- Blankenburgs G., J. Australian Inst. of Metals, 14/4 (1969) 236-241 8- Quenisset J. M., Fedou R., Girot F. A. and Lepeti tcorps Y., Proc.

Symp. on Advances in Cast reinforced Met. Composites, 88th World Mat. Congr., Chicago, Sept 25-30, 1988, pp. 133-138

162

9 with alumina film

150 -----""'"--without

film

(Laurent) 30 (Kohler)

600 700 800 900 1000

Fig 1 Contact angles between aluminum and a SiC plane substrate, as a function of temperature

1000

900 L

800 u . 755

o 20 6KF. 4AlF3

32.2 627

L + a-A1203 + K3A1F6

520

40 60 80 100

MOL (%) a-A1203

Fig. 2 Calculated "6KF.4AlF3" _ a-A1203 vertical section of the complex KF_AlF3-A1203-K20 system

500 1000 1500

a) as received b) K2ZrF6 (2.6.10 mg/cm2 )

~ annealing 650'C . -- c) Aluminum . w 0: d) K2ZrF6 3 (3 .6.11.25 mg/cm2 ) ~ Aluminum t:..

t:.. o ;... E­H ..J H to < c::l g c:::

d)

2500 3000

b) " ;

o 99 '... . i [j j 5 90

/11 i j .,: ill !

. : :'~ S 50 c);' n:" [

=/ ! a)! ::; ~ ::: g 20

• ;10 ci 0

' 00

10

'. 5

.. F:ULURE TENSILE STRESS (MP<I)

163

Fig, 3 Weibull plots for SiC CVD-filaments at different steps of the K2ZrF6 process (gauge length 40 mm)

Fig. 4 SEN micrograph of Al4C3 crystals formed at the surface of a SiC-CVD filament (after K?ZrF6 pretreatment and annealing 15 min in aluminum at 800 C) (x 10 000)

Impregnation untreated SiC-SiC SiC-SiC preforms parameters preforms treated with K2ZrF6

Zp max (c~) 1.5 ~ (no Al solidification) TZp max ( C) 660 784 PZp (MPa) 0.03 0.008 (for 25 em)

Table 1 Calculated impregnation parameters for SiC-SiC preforms with an overall porosity of 25%

ABSTRACT

MECHANISMS FOR MECHANICAL PROPERTY ENHANCEMENT OF FIBRE REINFORCED METALS

USING A HYBRID TECHNIQUE

EA FEEST, R.M.K. YOUNG, S.1. YAMADA", S.1. TOWATA*

Harwell Laboratory Metals Technology Centre, Building 35 - OX11 ORA DIDCOT - England

"Toyota C. R. D.L. - AICHI-KEN 480-11 - Japan

The properties of multifilament SiCIAI composites can be enhanced when they are 'hybridised' by small additions of particulate or whisker material. Previous studies have shown that a mechanism for this improvement is the reduction in the number of fibre contacts. The present work investigated other possible mechanisms. It showed evidence for the reduction of meniscus penetration defects and for the importance of the increased availability of the reinforcement/matrix interfacial area in reducing interface related degradation. In the AI-Cu matrix system this interfacial effect was manifested in the distribution of intermetallic in the composite.

1 - INTRODUCTION

Hybridisation by the addition of relatively small quantities of particulate or whisker material has been shown to enhance both the longitudinal and transverse strengths of unidirectionally reinforced metals Ill. In such hybridised composites the fibres are coated with fine SiC particulate or whisker material before the fibre tows are laid up into preforms and pressure infiltrated. This causes a struc­tural change in the composite since the adhering particles or whiskers prevent direct fibre contact. Thus melt penetration between the fibres is facilitated. Hybridised composites therefore have a better spatial distribution of fibres with fewer defects between them and much improved mechanical properties have been obtained. Enhancement of the transverse properties is particularly important in that many potential applications of unidirectionally reinforced metals are limited by the poor transverse properties so far achieved. The beneficial effects of

166

hybridisation have been shown to be alloy dependent which suggests that more subtle solute-related mechanisms may be relevant in addition to the above structural changes. Experiments on fibres extracted from the Al-Cu matrix system have suggested that the degree of fibre degradation in the unhybridised composite could not be explained by fibre attack alone. This points to the operation of a further alloy dependent degradation mechanism.

The aim of this study was to investigate the possibility that the property enhancement in alloyed matrix hybrid composites could be due to the influence of the markedly enhanced reinforcement/matrix interfacial area provided by hybridisation on these alloy dependent degradation mechanisms.

2 - EXPERIMENTAL PROCEDURES

2.1 Fabrication

Six specimens, designated A - F were supplied by Toyota CRDL. In all cases the continuous reinforcement was Tyranno fibre (-Ube). The matrix composition was either nominally pure Al or AI-4~ Cu, and the hybridising material was SiC particulate or whisker material. The particulate material was Betarundum <-Ibiden) with a mean particle diameter of 0.28 ~m,and the whisker was Tokamax (-Tokai Carbon) with lengths in the range 50-200 ~m and diameters in the range 0.2- 0.5 ~m. Composites were fabricated by a squeeze casting route, in which 19 * 150 * 2.2 mm fibre preforms contained in steel boxes were infiltrated at a pressure of 90 MFa maintained for one minute. Infiltration was performed with both preform and melt at 993 K and the specimens were cut out of the resulting blocks and ground to shape. The specimens were supplied as coupons for testing of longitudinal or transverse properties according to the designations given in Table 1.

2.2 Testing

The mechanical properties of the composites were evaluated by 3-point flexural testing <Toyota CRDL) and by longitudinal or transverse tensile testing (N.P.L., UK). For each type of speCimen, two separate castings were used for the tensile testing and four flexural test specimens were cut from a third casting. 8 mm wide, 1.8 mm deep flexure specimens were tested with a span of 50 mm and a ram speed of 1 mm/minute. The longitudinal and transverse tensile test geometries were 150 * 19 * 2.2 mm and 88 * 22* 2.8 mm respectively.

2.3 PhYSical examination

Fracture surfaces were cut from the tensile test specimens and examined using SEM. Metallographic sections cut from the specimens after tensile testing were polished and examined by optical microscopy and backscattered SEM. Fibres extracted from the composite by dissolving out the matrix in dilute NaOH were also examined using SEM. SpeCimens for TEM examination were prepared by' spark eroding 3mm diameter discs out of sections cut from the tab end of the tensile speCimens, dimple grinding to approximately 30 ~m and ion beam milling

167

until perforation. These were examined using a Phillips EM 400 TEM and a VG 501 cold source field emission STEM.

3 - RESULTS

3.1 Structures

In specimen A (non-hybrid), the fibres were touching with a mean number of contacts in a sample of 60 of 3.65. Macroscopic fibre-free bands some 5 or 6 fibre diameters wide existed through the specimen. TEM showed that the meniscus had failed to penetrate the contact zone between the fibres and that the defect thus formed is approxima'ely 1.5 ~m long compared with the fibre diameter of 10 - 15 ~m.

Intermetallic appeared to be associated with fibres but its appearance was irregular. In specimen B, addition of particulate to the material had reduced the coordination number to near 0, nearly all fibres appearing to be completely isolated from adjar.~nt fibres. The structure appeared to be more uniform than A but there were still large fibre-free bands which were possibly larger than those in A. In specimen C the whisker hybridised material was similar to B but there was less separation between the fibres, so that some fibres were actually touching. Whiskers were seen to concentrate in the fibre free zones. Larger fibre or whisker free zones were also evident. The structure of specimen 0 was similar to that of B except that the matrix was of pure aluminium instead of Al-4% Cu. Slightly more aggregation of the particulate was seen but fibre separation was good.

The structural observations were consistent with those reported in previous studies of hybrid composites, apart from the presence of macrodefects corresponding to fibre depeleted regions and the heavy aggregation of particulate or whisker. All the microstructural evidence demonstrated that there was an association of the inter­metallic with the particulate or whisker material, although this was not necessarily completely covered. This is clearly shown in a composition map taken with the STEM facility (Fig. 1). The 0.4 ~m diameter particle was approximately 80 % covered by intermetallic, and the adjacent region at the fibre matrix interface appeared to be depleted. The thickness of the intermetallic layer on both the particle and the fibre was variable, and up to approximately 0.08 ~m thick. No localised solute enrichment at the fibre/matrix interface was detected, nor was there any unambiguous evidence of fibre attack in the SEM investigation on deep etched specimens. TEM imaging of specimen C suggested that the scale of the intermetallic associated with the whiskers was larger, typically 0.2 ~m diameter.

3.2 Fracture surfaces

In specimen A there were flat regions where the fracture surface had propagated across the contact zone between adjacent fibres. No fibre pull-out was observed and there appeared to be little fracture debris on the surface. In specimens Band C, none of the flat regions visible in A were seen and the fracture surface was very rough with much debris in evidence. Little fibre pullout was observed. There was no evidence of direct crack propagation between adjacent (but not

168

touching) fibres. In specimen 0, some fibre pull-out had taken place. The fracture surface was highly irregular and individual fibre fracture events were well separated from each other. In specimen E the fracture surface was macroscopically planar and the fracture path tended to follow the fibre-matrix interface. The fracture surface in specimen F was inclined at 45' to the plane of the specimen indicating a shear failure. This appeared to coincide with a fibre depleted zone delineated at the edge of the specimen.

3.3 Mechanical properties

Properties from the testing programme are summarised in Table 1.

4 - DISCUSSION

4.1 Mechanical properties

The mechanical properties presented in Table 1 show a consistency in the ranking order of the strength of composites measured either by flexural or by tensile testing. The ratio of flexural to tensile strengths is however higher than has been observed in other related investigations Ill. Reasons for this apparent discrepancy are currently under investigation. One factor under consideration is the sensitivity of the tensile strength to the stressed volumes in differ­ent specimen configurations, particularly in view of the presence of macro-defects in the composite microstructures. However, the absolute values of the strengths do not affect the study of mechanisms reported herein. Key observations in this study are that the addition of 4% Cu to the aluminium matrix markedly reduces the longitudinal strength (samples B & D) and increases the transverse strength (E & F) and that hybridisation increases longitudinal strength (A, B & C).

4.2 Fibre distribution

Hybridising the composite has a marked effect on the fibre distribution. The poor fibre distribution in Specimen A comprising regions of high fibre packing density and substantial fibre depleted regions is attributed to attractive inter-fibre forces which are developed during infiltration. Hybridising the structure mechanically constrains the approach of adjacent fibres and also leads to a substantial decrease in the attractive forces between them 12/.

4.3 Meniscus defects

Unidirectional composites produced by infiltrating tows of fibre with liquid aluminium under pressure are prone to the formation of defects in the contact regions between fibres, since melt penetration in non-wetting systems, as for example in AI-SiC, is limited by the finite meniscus curvature that can be generated by the applied pressure 13/. Application of high pressures leads to improved melt penetration, but even at the pressures generated by squeeze casting (100 MPa), defects between 10 ~m fibres are still typically 1 ~m wide. The present stUdy showed that hybridisation reduced the

incidence of such defects (Fig. 2) because the narrow contact zone between touching fibres is eliminated.

4.4 Interface related effects

169

This study has indicated that hybridisation modifies the size and distribution of intermetailic second phase In the AI-Cu matrix (Figs. 3 a,b). For typical hybridised composites ~approx. 5 vol % hybrid addition) the calculated interfacial areas for whisker and particulate hybrids are respectively approximately 3 and 6 times those available in the unhybridised equivalent. This would lead to corresponding reductions in intermetallic layer thicknesses if they were uniformly distributed around the available Interfaces. In the hybrids, particul­ate material is normally located at the fibre matrix interface which is where solute-rich material is segregated during solidification, and this will strongly influence the resultant intermetallic morphologies. Whiskers might be less efficient than particulate at redistributing the intermetallic because much of the whisker is remote from the the fibre surface where the copper concentrations occur. The presence of the hybridising additions would also be expected to influence inter­metallic morphologies through nucleation, growth and coarsening effects. Refinement of the intermetallic phase distribution In the composite would enhance the strength of the composite and so it is likely that hybridlsation improves the mechanical properties as a result of increased interfacial area. This increase in interfacial area could also reduce chemical degradation of the fibre by reactive solutes. No evidence for this chemical degradation arose from the present microstructural studies. Such a mechanism would be expected to be more significant in more reactive systems such as SiC fibre /AI-Mg.

5 - CONCLUSIONS

Three possible mechanisms for property enhancement by hybridising the composite have been identified as follows:

Fibres are physically isolated from each other. Complete fibre separation in the particulate hybrid and partial separation in the whisker hybrid have been confirmed. Hybridising reduces the tendency to form meniscus penetration defects in the casting at fibre contacts. Hybridising increases the amount of interfacial area available. In the SiC/AI-Cu system this leads to a reduction in the scale of intermetallic formation at the fibre/matrix interface. In reactive systems the degree of fibre degradation is likely to be reduced

ACKNOWLEDGEMENTS

This work was carried out with the support of the Metal Matrix Composites Working Party and U. K. A. E. A. Underlying Research Program~le.

References 1. S. Towata and 2. R.M.K. Young, 3. A. Mortensen,

S. Yamada, Trans. Jap. Inst. Met, 27 9 (1986) P 709 AERE R- 13330, November 1988 Met. Trans. A ~ 1160

170

Table Specimen designations and mean properties of Tyranno tibre reinforced aluminium (bracketed numbers are best values)

Specimen Ma t11 x Hybrid Long, flexural Addl tlon strength (SPa)

A AH% Cu 0,42 (0,45) B Al -U Cu SIC .. 1,03 (1 ,06) C AI-4% Cu SICw 0,71 (0,78) 0 Al SIC p 1,4 (1,52) E AH~ Cu SICp/w F Al SlC p / w

Fig, 1 Alleu composition map In B taken with the STEM facility

Long, tens il e Transverse tensile strength (SPa) strength (GPa)

0,16 (0,18) 0,36 (0,37) 0,24 (0,28) 0,7:3 (0,79)

0,18 (0 ,19) 0,08 (0,08)

Fig 2, MeniSCUs penetration de f ec tin A l TEM)

Fig , 3 Size distribution of second phases in Specimens A and B (back-scattered SEM) 1 flm 1-1

POWDER METALLURGICAL PRODUCTION OF WHISKER REINFORCED MAGNESIUM

K.U. KAINER, J. SCHRODER, B.L. MORDIKE

Institut tilr Werkstoffkunde Technical University of Clausthal Agricolastrasse 6, 3392 CLAUSTHAL-ZELLERFELD - West Germany

ABSTRACT

The advantages of whisker reinforced magnesium are stiffnes,improved hardness, good wear resistance, reduction in thermal expansion and high strength:weight ratio. The disadvantages are expense of whiskers, high production costs of composites as well as health problems due to the carcogenic effects of fine whiskers. In this paper the preparation of whisker strengthened magnesium is demonstrated on the system magnesium-SiC or potassium-titanate whiskers. The properties of such composites eg modulus, strength, wear resistance, hardness and thermal expansion are discussed and compared with commercial Mg-SiC-F9-whisker composites (ARCO).

INTRODUCTION

Magnesium matrix composite materials offer several advantages when high performance low density materials are required. Conventional magnesium alloys possess several disadvantages eg low E-modulus, low high temperature strength and poor creep resistance, low hardness, poor wear resistance and high coefficient of thermal expansion. The aim in composite material technology is to eliminate these definen­cies. There are several techniques available to strengthen the mag­nesium matrix eg inclusion of non metallic inorganic short fibres or whiskers of high E-modulus, strength and hardness. The manufacture of short fibre or whisker reinfored magnesium composite materials can be accomplished by either infiltration of fibre preforms or by powder metallurgical methods. Since liquid magnesium is very reactive the powder methods offer certain advantages as there is virtually no or only slight reaction with the fibres or whiskers. The problems in the

172

powder metallurgical production are the mixing of short fibres and powder and the damage of short fibres during consolidation. Conse­quently little information is available on fibre or, particularly, whisker reinforced alloys. The results available show interesting property profiles (H.J. Rack and P.W. Niskame, ARCO Metals). In the following, the methods of manufacture and the influence of matrix and whisker on the properties will be described and compared with proper­ties of commercially available Mg-SiC-F9-whisker-composites.

I - MANUFACTURE OF COMPOSITE MATERIALS

Metal powder with the necessary size distribution is required if homo­geneous powder/whisker'mixtures are to be prepared. Magnesium powder (AZ91 and ZE63) produced'by the firm PSI, Polegate UK in a "Hermiga" gas atomizer was used. It is possible with this equipment to produce rapidly quenched magnesium powder w~th me~ grain size of ~ 15 ~ (Fig. la) for quenching rates of 15 - 10 K/s (W.G. Hopkins, K.U. Kainer et al). Table 1 shows the chemical composition of the alloys and the mean particle size. The SiC whiskers used were obtained from Tokai Carbon and the K20x6Ti02-whisker from Isolite Insulating Prod •• Table 2 lists the most important properties of the whiskers. Figs. Ib,c shows the morphology of the whiskers. The SiC-whiskers are finer (Fig. Ib) and show a higher aspect ratio than the potassium­titanate-whiskers (Fig. lc). The manufacture of the whisker reinfored magnesium composites followed the route described in Fig. 2. The most important steps in the production are the declustering of the whisker, the production of a excellent blend and the degassing of the blend. These steps influence the properties of the extruded materials.

II - MICROSTRUCTURE AND PROPERTIES OF THE MATERIALS

The microstructure after extrusion is similar for both whisker types. Fig. 3 shows a longitudinal section for the alloy AZ91+20 vol.% whisker (Fig.3a) in comparision with the microstructure of the commercial Mg-2Ovol.% SiC-F9-whisker-composite (Fig. 3b). Damage to the whisker is apparent in both cases which has led to a reduction in aspect ratio. The distribution of the whisker is more homogenous in the ARCO-material, although there are also some whisker-cluster. In Fig. 3c the whiskers are damaged and some oxide particles of the atomized powders are visible. the degree of whisker damage in the composites with atomized AZ 91 powder is higher than in the composites using low strength unalloyed magnesium as matrix alloy. In all cases adding whiskers leads to an increase in hardness, which in detail depends on the matrix alloy and whisker type (Fig. 4). The E-Modulus is also increased (Fig. 5). The best values were obtained for the ARCO material for which the the best distribution of whiskers and least whisker damage were attained. The yield stress (Fig. 6) of composites is very high when atomized powder is used. Unlike ARCO marterial a decrease in yield stress with increasing temperature is only observed above 100~C. At higher temperatures, however, there is a rapid fall in strength due to the inhomogeneous microstructure and the greater damage of the whiskers using atomized powder. Similar be­haviour is observed for the UTS (Fig. 7). Noticeable is the increase

173

in strength when atomized powder is used without a second phase. This is due to the fine grain size of the powder and the included oxide particles. Addition of whiskers to magnesium materials causes a reduc­tion in the coefficient of thermal expansion (Fig. 8). The reduction depends on the volume fraction of whiskers and their properties. The The wear properties are improved by addition of whiskers, the degree of improvement depending again on volume rraction and nature of the whiskers (Fig. 9).

CONCLUSIONS

The results show that an improvement can be made in the important properties of magnesium materials by the addition of whiskers. It is difficult to distribute the whiskers uniformly and without damage in the matrix. Without this optimum h~rdening cannot be achieved. Use of gas atomized magnesium powder enables the whiskers to be distributed uniformly but the high strength leads to a greater damage of the fibres. The strengthening is then comparable with the particle strengthening of magnesium.

REFERENCES

1. H. J. Rack and P. W. Niskame; Light Metal Age, (1984), 9 - 12. 2. ARCO Metals Comp., Greer, USA, brochure (1982). 3. W. G. Hopkins; Metal Powder Report, 42, (1987). 4. K. U. Kainer, W. G. Hopkins and B. L:-Mordike, to be published. 5. Tokai Carbon Co. ,Ltd, Tokyo, Japan, brochure (1986). 6. Isolite Insulating Products Co. Ltd, Osaka, Japan, brochure (1987).

chemical composition

mean particle size d50

AZ91

Al: 9.5% Zn: 0.4%

13.0 ].JlIl

ZE63

Zn: 6.0%, RE: 2.5%, Zr: 0.7%

14.1 ].JlIl

Table 1: Composition and particle size of gas atomized Mg-a110ys.

diameter (].JlIl) length (].JlIl) aspect ratio lId density (g/cc) tensile strength (GPa) tensile modulus (GPa) Mohs hardness

SiC (Tokai)

0.1 - 1.0 30 - 100 50 - 200

3.19 3 - 14

400 - 700 > 9

SiC (Silag F9)

0.3 - 1.3 50 - 500

150 - 1500 3.20

3 - 14 400 - 650

> 9

0.5 - 3.0 5 - 30

10 - 60 3.30 7

280 4

Table 2: Properties of the whiskers used (Tokai Carbon, ARCO Metals, Isolite Insulating Products).

174

Fig. 1 - SEM of powder and whiskers used: a) gas atomized AZ 91. b) SiC-whisker, c) K20x6Ti02-whisker.

rnagnlSium aUoy

I gas atomization

I mixing

I .nc'fSUlin9

deg.ssing

I hot pr.ssing

I

T I . hot deformation

I component component

whisk.r

I dlogglomtrotion

I

m.ln;1lQ I

component

Fig. 2 - Production route for magnesium­whiskers-composites.

Fig. 3 - Microstructure of extruded Magnesium matrix composites (longitudinal section): a) AZ91 + 20 vo1.% SiC-whisker, b) SiC-whisker and oxides removed from a). c) Mg + 20 vo1.% SiC-F9-whisker (ARCO).

,0

~

14 ." )

I 20

10

175

AZ91 ZE63 Mg -'11000 • "-!

I f j j i I I ~ t f ! ., 1 f ! j500 ! .. .. .

d ~ ~ ~ :r d

~ L

> t t ; 'I: ; i t ! Iil

1 ; ; i I J I! ~ ! e I ~ : : :. i ; a d d ; d ~

Fig. 4 - Vickers hardness as function of material composition.

AZ91 ZE63 M

-j ! f

·0

1 !

- ~ t Fig. 5 - E-Modulus of the Mg-

l ~ ~ whisker-composites as til til :r til til function of whisker-

f. f. I: i .. f. content, matrix compo-; J I J i I! ': . f ~ i ~ sition and whisker type.

t •• ting ,-,"11ft I'CI -_

• AZ" ahliDd pDVdIt

• AZ91.20val.%SiC ........

o 109 99.9511fG1111111 c Mg.20 ... % SiC-F9-........ •

1IIO

Fig. 6 - 0.2 proof stress of the materials depending on the composition and testing temperature.

176

' . .........................

• A1. 91 at.illd powdIr -.." ...............

100 -- ~ --_, --..... ~ -Al9t._%SiC-_

• 0 PIg 99.95 ......

-O~ .PIg._%SiC-F9-_"

-----~ O~ '~n o~:~_ ~ ..... ----~ ° ___ 0_

lOll

300 Inmv __ hn ("C I --

Fig. 7 - Ultimate tensile strength of the materials as function of the testing temperature.

AZ 91+SiC-whisker ZE63 Mg

Fig. 8 - Coefficient of thermal expansion.

.. '\. j

t i

.Al9I

.z"._.%SiC-_ oAZ"'-'%SiC-_ - -, 10 20 )0

ZE63

0", • 2Ovol.%SiC-whiIkIr

10 20 )0 10 10 )0

tilllllillin)--

Fig. 9 - Wear properties of the alloys and reinforced materials.

VIBRATORY ORIENTATION OF SHORT REINFORCING FIBERS IN METAL MATRIX COMPOSITES

ABSTRACT

B.SHPIGLER

Israel Institute of Metals-Technion R&D found Technion City - 32000 HAIFA - Israel

A vibratory treatment is proposed with a view to obtaining the controlled unidirectional orientation of short reinforcing cera.ic fibers in Metal Matrix Composites .anufactured with Powder Metallurgy methods. The treatment is applied to the metal-powder and fibers mixture before compaction and sintering, which results in the uni­directional orientation of the fibers that may either be uniform in the whole machine part or differ considerably in different zones. The effectiveness of the proposed mechania. has been proved over a wide range of experiaental para.eters, various vibratory systa.s and different powder fiber mixtures being used. The dyna.ic response of the mixtures as well as their dependence on frequency, acceleration, wave form, amplitude, and energy are discussed.

INTRODUCTION

Metal Matrix Composites (M.M.C.) combine the specific properties of a plastic metallic matrix with the high mechanical properties of a brittle reinforcing material (cera.ic, metallic, or organic fibers, whiskers, or particulates). M.M. Ca.posites reinforced with continuous unidirectional fibers show the highest strength [1]; but they are the most difficult and expensive to manufacture and are very sensitive to fiber da.age when additional .anufacturing technologies (forming, bending, extrusion, rolling) must be applied.

M.M. Ca.posites reinforced with discontinuous short fibers or whiskers are usually manufactured by one of three recently developed methods:

178

a. Fabrication of whisker or fiber prefo~ by pressure or sintering bonding followed by vacuu. or (high) pressure liquid Betal infiltration. In this case, the fiber size, orientation, and distribution depend ~stly on the prefon. ~rphology.

b. Ca.po-casting - which consists in the casting of previously .ixed and continuously stirred liquid (or sa.i-liquid) Belts of a .ixture of Betal and fibers. The orientation and distribution of the reinforcing fibera are isotropic and relatively non-unifona and a~st i~sible to control [2J.

c. Use of Powder Metallurgy (PM) technology, consisting in blending a .ixture of short fibers and Betal powder followed by co.paction and densification by (cold) pressing followed, in turn, by sintering, usually below the solidus te.perature of the Betal .trix.

All three Bethods yield shapes or bi llets having an isotropic mor­pholo~y with randomly oriented and distributed fib~rs and considerably lower ~hanical properties (IJ. To achieve the unidirectional arrang~t of the fibers, additional .anufacturing technologies (such as extrusion, forging or rolling) are introduced which, however, cause significant d .... e to the fibers, considerably reducing their average length (I) and shape factor (lid) and therefore their reinforcing effect (2J.

A recently reported coaparison of Bechanica} properties of SiC fiber reinforced alu.inu. .trix Ca.posites fabricated by the processes Bentioned showed a notable advantage of the PM method [3J, .ainly due to the ha.ogeneous distribution of the SiC reinforcement in the .atrix. Accordingly, it is widely reco .. ended [4] that the main target of future R&D work be: a. To increase the fiber aspect ratio (reduce damage) b. To ensure proper fiber alignment.

I - THBORBTICAL APPROACH TO FIBER - MATRIX LOAD TRANSFER

The contribution of the fibers to the mechanical properties of the M.M. Ca.posite depends on their quantity (volu.e fraction Vf ), orientation (80), length (1) and aspect ratio lid) [2].

Kelly and Davis (5J have shown that the composite strength ('c) with fully aligned (unidirectional) reinforcement is given by the equation:

where

'c = 'fvfCI - lc) + '.Cl-vf )

21 'f = Fiber strength

I = Actual fiber length I = The critical fiber length - defined as: c

lc = df 'f

'y df = Fiber's diameter

, = Flow stress in .. trix y

Nick et al. (6] have investigated the microstructure of PM

179

2124/15' SiC M.M.C. after extrusion. They have found that the w

preferred orientation of the whiskers was far fra. perfect. while their average aspect ratio. lid. had critically decreased fra. 50 in the as sintered condition to =4 after extrusion. as a result of the dllll&ge caused by the extrusion operation. The siCOificance of this apparently inevitable ~e is that the reinforcing ele.eots .eke a considerably lower contribution to the ca.poeite strength and stiffness than could be expected.

Theoretical predictions of the effective teoaile .adulus aade by Rack. Baruch. and Cook [7) .how a hither contribution to the asterial's stiffness by the higher values of the aspect ratio of the reinforCell8Dt.

Most exper~tal results available on the dependence of M.M. Ca.posite .trength on fiber orientation were obtained for continuous fiber reinforce.eot [8).

The above data and other. that are available show that the dependence of M.M. Ca.poeites properties on fiber orientation i. very strong but probably .are ca.plicated than the .anotonic function predicted by.a.e inv .. tigators [9). Neverthel.... there i. no doubt that M.M.C. properties strongly depend on the orientation and length of the short fibers. Finding a better way to control the orientation of the reinforcing fibers while preventing significant fiber dllll&ge can greatly ~rove the fiber.' contribution to the ca.poaite structure and properties.

II - THE BXPERIMENTAL APPROACH

During the last few years. efforts have been .. de by various researchers to develop .ethods to gain control over fiber orientation in M.M.C. Allan and Bevis [10) have reported the develo.-mt. at Bruael University. G.B •• of a ~ltiple live-feed injection .aIding syst~" (MLFM) capable of controlling the orientation of fibers in thel1lOplastic cOJlPOSitee aade by (liquid) injection .aldin,. The method is based on the CCIIbined application of pressure pul... to two separate feeding lines of .alten .. terial. The shear actin, on the .elt before and during solidification i~roves both the packing and the orientation distribution of the fibers as well as the ca.posite mechanical properties [10).

Masuda and Itoh [11) of the University of Tokyo and I.C.I. [12) announced in 1987. the develo~nt of an "Aligned All.-ina Fibre" pre fora asde fra. short A120, fibers suspended in freon. They are electrostatically aligned. then sedi.uted to fora a prefora (II8t). and finally dried. These prefor.. were infiltrated with liquid AI. alloy which for.. a unidirectional co.posite .arpbology with very little fiber dllll&ge and ~roved .echanical properties. No work is known to have been done using vibration to obtain controlled anisotropy in powder .ixtures or in M.M. Ca.poeites aanufactured by Powder Metallurgy technology.

Our method to control the orientation of reinforcing fibers in Composi tes made by the Powder Metallurgy Technology is based on the beneficial effects of vibratory treatJIIeDt. which has already proved effective in the ca.paction and densification of granular beds in general and powders in particular [13. 14.15].

180

The vibratory techniques concerned involve the supply of energy increments to the bed of particles at a specific frequency for a selected time. This technique has been under evaluation by our laboratory during the last years (15,17]. The effectiveness of the vibratory treatments was confirmed and found to depend on several major parameters: a. Powder (mixture) characteristics: Shape and size distribution,

volumetric ratios, friction coefficients, surface roughness, hardness, density, etc.

b. Vibration parameters: Frequency (F), amplitude (A), acceleration (g) direction (v), wave form and energy, etc.

The experimental set included 3 and different systems.

2.1 A low-frequency mechanical motor-driven vibratory system

This comprised a specially designed multi lever arms system activated eccentrically by an electric motor (a.c.)through a 250-2500 rpm continuous gear (see Fig. 1). The system can operate in various conditions, continuously and independently changing the frequency range (4 to 40Hz), the acceleration (0 to 100g), and the amplitude range (0 to 4Omm) using various wave forms. The powder-fibers mixture specimen is kept in a transparent vessel and is mounted at the top of the vertical arm, so that it can be observed during the process with the aid of a synchronized strobe light.

2.2 A pneumatically activated high-frequency vibratory system

This consisted of a flat holding plate connected to (but dynamically isolated from) a holding frame through a variable number of soft coil springs (see Fig. 2).

The vibrational movement was obtained using with types of pneumatic vibrators operating at a number of frequencies, amplitudes, accelerations, and wave forms, some of these parameters strongly interdependent. They were measured by means of accelerometers mounted at different locations on the plate and X-V oscilloscope display.

The powder-fibers mixture specimens in their transparent containers were mounted at different locations on the plate, the movements of the mixture being observed through the transparent wall with or without stroboscopic illumination.

2.3 A small-size model vibrational system

This device consisted of a smaller holding plate adapted to the use of special cans filled with a variety of differently coloured and insoluble liquids (see Fig. 3). The behaviour of this continuous non-particulate fluid was observed during the different vibratory treatments and could be used as an indication of the expected behaviour of the powder-fibers mixture and its dependence on the treatment parameters.

The specimens used at this stage of the research consisted of "Acryfix", a non-metallic white powder, in mixture with black SiC or metal fibers. At the end of the experimental vibratory treatment the air in the can was slowly evacuated to a pressure of approx. 5-10 mTorr and an adequate amount of liquid hardener infiltrated in order

181

to obtain a transparent solid sa.ple of the ca.poeite, in which the fibers distribution and orientation can be clearly seen (see Fig. 4). A thorough exa.ination of the effects 'of these operations on the mixture aorpbology proved that they do not in any way affect the fiber distribution or orientation.

The evaluation of the treat.ent's effeCtiveness included: a. In-situ visual observation of the .ixture's reactions and of the

fiber aligoaent (see Fig. 4). b. Microstructural examination by optical .icroscope. c. Quantitative Image Analysis of fibers length and orientation.

BXPBRIMBNTAL RESULTS

1. Positive results were obtained with both aotor-driven low-frequency, and pneumatically driven high frequency, vibrational syste.s.

2. The required wave forms and a.plitudes required differed considerably from one frequency to the other. MIlch higher allPlitudes and accelerators were required at the lower frequen­cies (f=lO+4O Hz) than in the high frequency range (f > 150 Hz).

3. When it was successful, the vibratory treatment reduced the orientation distribution range around a controllable central symmetry axis to an opttaua of * 100 •

4. The fiber orientation distribution achieved can either be uniform in the whole can or differ considerably in various regions of the .ixture, depending on the parameters of the vibratory treatments and the dyna.ic properties of the container and holding device (size and shape, stiffness and elasticity, natural (resonance) frequency, etc.).

5. For each specific situation (.ixture, vessel and vibratory parameters) a different period of time was required for best results, with satisfactory repeatability. Continuing the treatment beyond that period was liable to cause considerable fiber deaage without further improving the orientation distribution and sometimes worsening it.

CONCLUSIONS

1. Vibratory treatments combined with powder metallurgy technology appear to offer a promising and reliable new method for the aanufacture of ca.poeite materials reinforced by short fibers marked by controlled anisotropy and improved properties.

2. The proposed method can be applied to both Metal and Cerllllic Matrix Composites.

3. The addition of vibratory treatments to P/M manufacturing of composite materials may lead to many evident advantages: a. Avoidance of additional aanufacturing processes - such as

extrusion, forging, or rolling. b. Reduction or eli.ination of fiber damage and improved

properties. c. Direct aanufacture of final shape or net-shape M.M.C.

products. d. Improved distribution of the reinforcement.

4. The application of this method to full-scale manufacturing

182

processes still requires ample R&D work for the detenaination of vibratory treatment paraaeters as well as for the development of adequate and adaptable vibratory systems.

REFERENCES

1. Mel M. Schwartz - "COIIPOsite Materials Handbook", McGraw-Hill Book Company, New York, N •. 1983, p. 237.

2. Christophe Milliere, "Elaboration, Structure et Proprietes de Materiaux COIIPOsites a Matrice Metallique", Ph.D. Thesis, Institut National Polytechnique de Grenoble, France, 6 May 1986 -p. 186.

3. F. Klausen and R.F. Singer, "Influence of Processing on the Mechanical properties of SiC Whisker Reinforced Aluminum Compo­sites", 31st Intern. SAMPB Symposium April 7-10, 1986, p. 1701.

4. J. Dinwoodie and I. Horsfall, "New Developments with short staple Alumina Fibers in Metal Matrix Composites", Proceedings, ICCM VI & ECCM II, London 1987, Vol. 2, p. 2.394-2.395.

5. A. Kelly and G.J. Davies, "The principles of the Fiber Reinforcement of Metals", Met. Reviews 10(1965), p. 1.

6. T.G. Nieh, R.A. Rainene and D.J. Chellman "Microstructure and Fracture in SiC Whisker Reinforced 2124 Aluminum Composite", Proc. ICaM-5, ed. by W.C. Harrigan et al., San Diego, 1985, p. 825.

7. H.J. Rack, T.R. Baruch and J.L. Cook, "Mechanical Behavior of Silicon Carbide Whiskers Reinforced Aluminum Alloy", Progress in Science and Engineering, Proc. of ICaM-IV, Tokyo, 1982, eds. Kswata et al., p. 1465.

8. T.W. Chou, A. Kelly and A. Okura, "Fibre-Reinforced Metal-Matrix Composites", Composites, Vol. 16 No.3, 1985, p. 187.

9. M.M. Schwartz, "Composite Materials Handbook", by McGraw-Hill, 1984, p. 315.

10. P.S. Allan and M.J. Bevis, "Multiple-live-feed injection moulding", Plastics and Rubber Processing and Applications, Vol. 7, No.1, 1987, p. 3-10.

11. Senichi Masuda and Tomohito !toh, "Electrostatic means for fabrication of fiber-reinforced metals", (to be) presented at the IEEE/IAS 1987 Annual Conference (Oct. 1987) Atlanta, Georgia, U.S.A. (5 p.).

12. I.C. I. "SAFIMAX-Aligned Alumina Fiber" - An advanced fibre from I.C.I. - commercial brochure (4. p.).

13. H.H. Hausner, Kempton H. Roll, P.K. Johnson, "Vibratory Compact­ing, Principles and Methods", Plenum Press, New York, 1967.

14. R.T. Dotter, "Blending and Premixing of Metal Powders" A.S.M. Metals Handbook Vol. 7; 9th Edition - "Powder Metallurgy" pp. 186-189; 306.

15. Henry H. Haisner, "Handbook of Powder Metallurgy", Chemical Publishing Co., Inc. New York, N.Y., 1973, pp. 118-121.

16. B. Shpigler, U. Betsalel, o. Botstein, "Metal Matrix Composites -by Powder Metallurgy" Res. RPt. 2-5045-59,' lsr. lnst. of Metals, Technion, July 1986, pp. 34-61.

17. B. Shpigler, o. Botstein, "Fiber Reinforced Sintered Heavy Metal, RPt. Stage I, Res •. Rpt. 5045-68, Isr. Inst. of Metals, Technion Oct. 1986. ~

Fig.1

Fig.)

Low frequency - Motor driven Fig. 2 vibratory system 1. Motor 4. Powder can 2. Eccenter 5. Acceleraneters 3. MIll tilever measurement

arms system system

Vibrational modeling system 1. Pneumatic 3. Liquid filled

vibrator can 2.Vibrating plate 4. Accelerometers

,., .. ,

183

High frequency - pneumatic vibra­tory syst ..

1. Specimens holding vibratin; plate 2. Vibrator and acceleraaeters 3. Meuur-.tts and display system 4. Powder cans

Fig.4 Fibers orientation obtained through vibratory treatment (transparent specimen)

Col

Determination of chopped fibe~s o~ientation in

powde~/fibe~8 mixtu~e by vib~ato~ t~eatmant

a. Initial i&ot~opic mixtu~e

b. Intermediate .taqe

c. rirral ataqe -unidirectional fibers o~ientation

TITANIUM MATRIX COMPOSITES REINFORCED BY C.V.D. FILAMENTS: A REVIEW OF THEIR

THERMO-MECHANICAL CAPABILITIES

Y. LE PETIT-CORPS, T. MACKE', R. PAILLER, J.M. QUENISSET'

ABSTRACT

Laboratoire de Chimie du Solide • Laboratoire de Genie Mecanique

Universite de ecrdeaux I 351 cours de la Liberation - 33405 TALENCE - France

A review of titanium matrix processing techniques is presented. Among the C.V.D. reinforcements commercially available, those coated with a thick protective layer (e.g. SCS-6 from AVCO or B4C/B from SNPE) are the only ones which strengthen effectively the titanium matrices. A synthesis of the chemical and mechanical (static and dynamic) behaviours of 1 D-SCS-6 I Ti-6AI-4V and 1 D-B4C/B I Ti-6AI-4V elaborated by vacuum hot pressing (V.H.P.) is given.

INTRODUCTION

As monolothic materials (e.g. titanium or super alloys) are used today to their limits in structural aerospace or aeronautical applications, new materials like metal matrix composites (M.M.C.) could be an alternative in the future for light structures which have to withstand high degrees of thermo mechanical sollicitations.

The potential of the refractory matrix composites reinforced by ceramic fibers are (i) structural components (panels, tubes, missile bodies or stabilizers), (ii) turbine engine components (fans, compressor blades).

Titanium or intermetallics (e.g. Ni3 AI) matrices reinforced by continuous or short ceramic fibers could be potential materials aimed at this purpose.

Though these two kinds of matrices are going to be in competition in the future, the related M.M.C's exhibit similar

186

features as to their processing and properties 11,21. Within the frame of the present contribution, we have decided to focus on titanium matrix composites (T.M.C.) reinforced by boron or silicon carbide C.V.D. filaments.

The interest of the titanium matrix lies in its high specific properties and its resistance to corrosion which account for its wide use in aeronautics for a long time.

The ceramic reinforcements have to improve the stress to failure, the fracture toughness, creep, fatigue properties as well as stiffness of the matrix.

A synthesis of the researches carried out on these composites for the last decade is proposed stressing the problems linked to the elaboration, the compatibility between the constituents as well as the mechanical properties.

1- PROCESSING TECHNIQUES

The processing difficulties of T.M.C's which have been reported by the first researchers are related to the high melting point. and the high degree of reactivity of the matrix. However, titanium and some of its alloys exhibit in a rather low temperature range (Le., below 10000 C) a superplasticity behaviour which has been turned into account in the elaboration of the composites.

1.1. Processing based on superplastic deformation

During the processing (V.H.P. or Rapid Omnidirectional Compaction (R.O.C.)), the C.V.D. filaments are aligned and spaced between titanium foils, then the preform is submitted to a temperature, pressure and time cycling step allowing (i) the plastic deformation of the matrix around the fibers, (ii) the chemical diffusion between the components 13,4/. Nowadays, these techniques are no more limited to simple shapes (tubes, panels) besides the materials can be used to make various complicated structural configurations by super plastic forming diffusion bonding (SPF/DB) 15/.

1.2. Processing based on powder matrices

In plasma deposition, a torch is required to spray metal droplets on the fibers 16/. During processing, in spite of a very short contact time between a nearly liquid matrix and the fibers (few millise-

187

conds), an interphase appears. The metal coating is however thin (few microns) and a compaction by V.H.P. is then necessary to achieve a complete material.

The powder metallurgy (P.M.) technique enables to process composites strengthened by short reinforcements with various chemical compositions (alumina fibers, whiskers, particulates ... ) 12/. One can expect from this technique to obtain complicated shapes with isotropiC properties.

2. CHEMICAL ANALYSIS OF THE FIM INTERFACE

During processing, a reaction zone is formed as mentioned above, at the FIM interface. Its thickness depends on the processing technique. (- 0.5 11m for V.H.P., _ 0.2 11m for plasma or P.M.) 12,3,61 The growth of this zone is then linked first to the nature of the constituents (SiC, B, C, Ti, AI ... ) .and to the processing parameters (Temperature : 750 - 9500 C), (Duration : few minutes - few hours).

The reaction growth rate can be controlled by modifying the chemical composition of the components (surface coatings on the fibers, alloying elements in the matrix).

The first filaments which have received C.V.D. surface coatings were boron fibers. Currently, available filaments are SiC based. The surface deposits thought after to protect the reinforcements and to avoid the FIM chemical reaction are (i) carbides or carbon layers (TiC, HfC, C), (ii) oxides (Zr02, Y203, Hf02 ) obtained by C.V.D. or P.V.D. 16/.

As an example, the SCS-6 fiber supplied by AVeO has a four micron carbon rich outer layer 171. When titanium reacts with this coating, titanium carbide is formed in a first step. As long as the fiber coating is not entirely consumed, the titanium carbide plays the role of a diffusion barrier by preventing the titanium from reacting with SiC. For long annealing durations, the released silicon can none the less diffuse towards the matrix giving rise to titanium silicide (Fig. 1). The reaction growth rate is then determined from the titanium and silicon kinetics through· the titanium carbide layer.

Vanadium, molybdenum and aluminum were found to cause marked reduction in reaction kinetics with silicon carbide fibers 18/.

In Ti-6AI-4V based composites, the alloying elements are pushed at the interface (Fig. 1). Vanadium does not generate stochiometric compounds at the interface whereas aluminum reacts with titanium giving rise to intermetallic products (e.g. Ti3AI) which play the role of a diffusion barrier and avoid the extension of titanium silicide 181

188

The development of high resolution chemical techniques (E.P.M.A, A.E.S., S.T.E.M., T.E.M .... ) led to determine accurately the chemical composition of the interphases in such systems 13, 9,10 I(Table I).

3. MECHANICAL PROPERTIES

Single fiber composites have been elaborated in order to predict (1) the chemical compatibility between fiber and matrix, (2) the interfacial shear strength.

For uncoated filaments (e.g. SiC(l:), SiC 4, B/w) a very short annealing time in presence of titanium results in a drastic decrease in U.T.S (Fig. 2). Titanium carbide or titanium diboride acts as stress concentrators on the filament surface. On the opposite, the U.T.S. of filaments with a thick protective coating (e.g. SCS-6, B4C/B) remains unchanged even after long annealing treatments (Fig. 2). Although reaction products were formed, the main C.V.D. part of the filament is still protected against notch effects due to microcracks taking place in the brittle reaction products which fail first(Fig. 3).

The Fraser mechanical test was applied to determine the interfacial bonding in single fiber composites elaborated by V.H.P. 111 I.

These results suggest that uncoated filaments exhibit a very high interfacial bonding (750 MPa for SiC 4 and 345 MPa for SiC (l:)). on the opposite the F/M adhesion does not seem to have reached its maximum value for the other filaments (240 MPa for B4 C/B and 180 MPa for SCS-6). The protective coating limits the fiber-matrix interdiffusion phenomena and enhances crack deviation along the fibers.

According to these results, 1 D-SCS-6 I Ti-6AI-4V and 1 D-B4C/B I Ti-6AI-4V materials with a 35 % volume fraction of fiber were elaborated by V.H.P. and then mechanically tested.

The improvement in the quality of the surface treatments of the fibers produced currently nowadays enables to obtain composites whose room temperature axial U.T.S. obeys the R.OM. (Fig. 4). The buffer coating improves the ductility of the composites (eR '" 1.2%).

Transverse properties are still rather poor (Fig. 4). This feature is due to a weak adhesion between the interphase coating and the silicon carbide substrate.

The interest of T.M.C's lies in their high temperature tensile properties . Unfortunately, only few values are available and more high temperature characterizations need to be done to extend the T.M.C's field.

189

In order to complete the characterization, dynamic tests (impact resistance, crack propagation and thermal cycling) were performed and gave the following results:

(1) The dynamic toughness measured by an instrumented Charpy apparatus is improved, compared to the unreinforced matrix (KIQd = 1 05 MPa~m (1 D-SCS-6/Ti-6A1-4V) 170 MPa~m (Ti-6AI-4V)). However the rupture work is significantly decreased (105 kJ/m2 for the T.M.C. and 300 kJ/m2 for the matrix. The brittle ceramic fibers lead to a decrease in the rupture time and, as a result, in rupture energy. An improvement in KIQd was observed for a 30 to 40 hours diffusion annealing at 850° C. This feature was related to a better F/M bonding 112/.

(2) A Ti-6AI-4V matrix material reinforced with B~C/B or SCS-6 fibers exhibited better fatigue life times (N = 107, N = 2.105) than the unreinforced matrix (N = 6.104) cycles to failure respectively for R = 0.1 and amax= 515 MPa. The longest fatigue life time was observed for B4C IB composites 14/.

Diffusion treatments at 850°C resulted in an increase in toughness and a decrease in crack growth rate when the reaction time was sufficient to promote a better F/M adhesion and a larger energy dissipating mechanism in the matrix around the splitting zone 113/.

(3) Thermal stresses induced by the C.T.E. mismatch result in thinner reaction zone due to voids formed around the filaments during thermal cycling. The load transfer between matrix and fibers is hardly achieved by this debonding at the F/M interface, the stress to failure is lowered, however the rupture work is increased. Whatever the chemical composition of the filament may be, the higher the cycling frequency is, the stronger the damage is 112/.

4 - CONCLUSIONS - PROSPECTS

The recent improvements which have been achieved regarding the basic materials (e.g. new C.V.D. fibers) and the research techniques (chemical or mechanical) contribute to a better quality of the T.M.C's.

Some chemical and mechanical data for SCS-6/Ti-6AI-4V and B4C/B/Ti-6AI-4V composites were given. These materials are likely to find soon applications in aerospace industry.

However, there is still a lack of reliability in the mechanical properties probably related to processing . It would be advisable to develop more extensively research on T.M.C's as it is presently done in the field of aluminum matrix composites.

190

REFERENCES

1 - A. Bose, B. Moore, N.S. Stoloff and RN. German, private communication (Nov 1987)

2 - G. Das and F.H. Froes, to be published by Les Editions de Physiques in Titanium, Science Technology and Applications, (Cannes June 1988)

3 - P. Martineau, R Pailler, M. Lahaye and R. Naslain, J. of Mater. Sci., 19 (1984) 2749

4 - Smith and F.H. Froes, J. of Metals, (1984) 36 (3) 19-26 5 - D.J. Chang and W.H. Kao, SAMPE Journal (March/April 1988)

13 -17 6 - R.R Kieschke and T.w. Clyne, to be published by les Editions

de Physiques in Titanium, Science Technology and Applications Cannes June 1988)

7 - Y. Le Petitcorps, M. Lahaye, R. Pailler and R Naslain, Composites Science and Technology 32 (1988), 31-35

8 - W.D. Brewer and J. Unnam, The Metallurgical SOciety of AIME, Conference Proceedings edited by J.E. Hack and M.F. Amateau (1983) 39-50

9 - G. Das And RE. Omlor, Proceedings of the 46th Annual Meeting of the electron Microscopy of America (San Francisco, 1988) 738-739

10 - H.J. Dudek, L.A. Larson and R. Browning, Surface and Interface Analysis, Vol. 6 N° 6, 1984 274-278

11 - Y. Le Petitcorps, R Pailler and R Naslain to be published in Composite Science and Technology, Special isue for M.M.C.'s

12 -K. Nakano, L. Albingre, R Pailler and J.M. Quenisset, J. Mat. Sc. Letters 4 (8) (1985) 1046-1050

13 - J.M. Quenisset, P. Soumelidis and R. Naslain, Proc. E.C.C.M. (A.R. Bunsell, P. Lamicq, A. Massiah eds), (Bordeaux 1985) 571-576

Products E (GPa) cr,MPa) ER(%)

TiB 2 540 1350 0.25

TiSi 2 265 1190 0.45

TiC 450 1350 0.30

Ti 5Si3 235 1500 0.64

Table I: Chemical species and mechanical properties of reaction zone products formed at the interface between C.V.D. filaments and titanium matrices.

191

~~--------------~~ SI

:;60 Ii

140

20

4 8 12,16 20 24 D'-tance,11"1

'".!!. co ....

An .... I ... lima. IIln

l' • .. z_ (I)

a) 00

00

000

Ln (I) Fig .1. Auger line scans on a cross. section of (a) SIC (I) fUaml'" from 1111"'" Composlle (FRG)

a SCS-6ITi-6AI-4V composite annealed (b) I fIIamanl from INP! (f)

140 hours at 950·C. (c) SIC 4 filament from AVCO (USA) (d) ICI-2 fUamant from AVCO (USA) (e) 14C11 filament from INPE (f) (I) sew nll_ from AVCO (USA)

l. :I

g co

I

Fig.2. UTS versus time for C.V.D. filaments annealed at 850·C with a titanium coating (1-2 ILIT1 thick).

100

Ii 80

60 ~J:

~ 8 40 18 20

0 0

2000

100 200 300 Spultarlng nnw, min

1500

Annaallng durallon al I5O'C • Hour

SCI .. I TI .. AI-4V

Axial

• Tran.v.r ••

Unrllnlorcld malrll

:. :I

1000 Fig.3. A.E.S. concentration profiles near the surface

i -------------------------------- i of a SCS-6 filament annealed at 850·C J!

I in presence of titanium. 500 I-

Fig.4. Stress to failure versus reaction zone thickness for a SCS-6/Ti-6AI-4V composite (Vf=35%). tested

in tension at room temperature.

ABSTRACT

A COMPARATIVE STUDY OF THE MECHANICAL BEHAVIOUR OF ZINC REINFORCED BY

STAINLESS STEEL FILAMENTS MANUFACTURED VIA TWO DIFFERENT PROCESSES

A. MADRONERO, M. PRENSA MARTINEZ-SANTOS

CENIM Avda. Gregorio del Amo, 8 - 28040 MADRID - Spain

Little attention was paid to zinc matrix composites in the past, due to the high density of zinc, but its very low melting tempera­ture allows a very easy manufacturing process and thus its indus­trial use as a structural material may be expected in the future.

In the case of zinc matrix reinforced by stainless steel fila­ments, a comparative study is performed in order to establish a relationship between manufacturing process and the mechanical features of this material.

In this paper, samples are manufactured by hot press under vacuum of SS/Zn monolayers. The matrix included in such monolayers was apported via electrolytic deposition or by liquid metal infil­tration. Just before incorporation in the zinc matrix composite, the stainless steel filament could optionally be surfacially activated by pickling in a bath having controlled acidity.

The mechanical performance of SS/Zn composites is strongly in­fluenced by the duration of the sintering period and activation of the steel filament. The fracture energy depends on the failure mode (transfibrilar or flexure breakdown) and surface pre-treatment of filaments. For electrodeposited matrix, short sintering periods yield better toughness, but for infiltrated matrix three hours consolidation are recommended.

INTRODUCTION Zinc matrix composites with 18/8 stainless steel filament as re­

inforcement, is a cheap composite for use in cases where no high specific strength is required. Its low process temperature suggests a possible use in high work of fracture applications. For example,

194

the sintering pressure for aluminium is about 70 MPa, while for zinc a figure of 40 MPa works quite well.

The procedures to apport the zinc matrix onto the steel fila­ments before sintering it by hot pressing under vacuum, are electro­lytic deposition (1) and infiltration by passing filaments through molten metal (2).

The surface state of reinforcer filament plays a leading role in this type of composites. The supply state of surface stainless steel filament consists of a continuous oxide layer having superimposed oxide platelets. This complex coating can be removed by a short pickling through a dilute sulphuric acid water solution, just before the zinc matrix is apported.

By changing the acid concentration we could establish that in concentrated solutions of less than 30'\ vol there was no damage to the continuous layer, but the oxide platelets were totally removed during a brief immersion. Also, for concentrations above 75'\ acid, aside from the clear removal of the thin oxide layer, pits appeared on the filament surface with consequent fiber strength damage. Therefore in this paper, concentrations of 30'\ and 75'10 are used as the two severity levels chosen for the activation levels of the acid pickling.

During the hot pressing stage a remarkable diffusion of zinc took place towards the filament outer layer, yielding a thick reac­tion zone (3) similar to other materials in which overstrengthening was observed (4) at higher temperatures.

According to our measurements of zinc concentration profiles observed by SEM, and using the ordinary calculus of Fick' s Law for cylindrical symmetry (5), we could establish a diffusion coefficient value D = 1,44.10-12 cm2s-1 when the reinforced 18/8 steel filament was in supply state, only cleaned with organic solvents. Measuring on an activated reinforcer filament after pickling in 75'\ acid solution, the value was D 6.74.l0-l4cm2s-l • Performing measurements changing temperature and diffusion period, it was pos­sible to calculate for the frequency factor the value Do = 0.63 cm2 s-l and for activation energy Q = 28.8 kcal/mol.

These figures explain the reaction zone structures observed dur­ing the manufacturing processes used in this paper (Fig.l and Fig.2). Using electrolytic deposition (room temperature) for zinc without pickling, no reactio~ zone was observed because oxide coat­ing acts as a diffusion barrier (Fig.l). As the molten zinc wets the steel filament to higher temperatures, the reaction zone is very significant, with immediate formation (Fig.2.b) if the surface fila­ment was activated; if after that diffusion is continued, the reac­tion zone becomes dissolved (Fig.2.c) and the softening phenomena appear (6).

The influence of these parameters, state of filament surface and manner of matrix apporting, are the main set of controllers of composite performances. The other set of controllers are the para­meters of the sintering process: sintering temperature (only a very few degrees below melting point) and pressure, versus time. There are two main ways to carry out the sintering process; a deep, cold compactation in atmosphere, followed by a short hot pressing under

195

vacuum, and a moderate slope heating under vacuum and constant pressure.

Both methods give an over increase of elastic modulus of the SS/Zn composite, and a diminishing of ultimate tensile stress, according to calculated theoretical values (6). The differences between the two procedures diminishes when the surface filament is very activated (acid pickling 75'\) and increases when the filament keeps its initial oxide coating (7).

TENSILE PARAMETERS By performing mechanical tests, it was possible to evaluate the

kinetic of the reaction zone displayed in Fig.1. According to the left plot of Fig.3, just after the second hour of the sintering process, the increased rigidity and reduction of strain failure, point to the role of the mentioned brittle interface.

The significance of whether or not to use acid pickling of the filaments, can be observed in the Table, where ~ and S constants are the adjustment coefficients of the rule of mixtures for elastic modulus and ultimate tensile stress respectively.

According to the right plot of Fig.3, the matrix apporting pro­cedure has only a moderate influence, in spite of the differences in matrix recrystalization and monolayer face welding during sintering time. At the end of this sintering period, the theoretical figures are raised.

FRACTURE ENERGY This aspect depends very strongly on interface strength, as

Fig.4 shows, and the failure mode according to the type of failure test. Iosipescu' s test reflects (8) transfibrilar failure for uni­directional composits. According to Fig.4, high fracture work is achieved with pickling in 75'\ acid and for short time sintering (infiltrated matrix) or long time sintering (electrodeposited matrix). Figures for fracture work larger than 300 kJm- 2 can be achieved.

APPLICATIONS The indus.trial application for this kind of SS/Zn material falls

within the area of low cost and high fracture energy. Fig.5 shows the difference in piercing of this material in com­

parison with steel. Besides a high value of energy consumption, the spreading of the damaged zone converts to SS/Zn in a candidate material for the rear layer material for a two layered plate to withstand a threat of piercing.

BIBLIOGRAPHY

(1) Harris, S.J.; Baker, A.A.; Hall, A.F. and Bache, R.J. Transactions Inst. Met. Finishing, 49 (1971) 205-215.

(2) Renton, W.J. (editor) "Hibrid and select metal-matrix composites", American Inst. Aeronautics and Astronautics, (1977) 166.

(3) Madronero, A. Rev. de Metalurgia, 21 (1985) 346-356.

196

(4) Petrasek, D.W. and Signorelly, R.A. "Tungsten alloy fiber reinforced nickel base allow composites for high temperature turbojet engine applications", ASTM Special Technical Publication, 460 (1969) 405-416.

(5) Darken, L.S. and Gurry, R.W. "Physical Chemistry of Metals", McGraw-Hill Book Co., (1953) 437-452.

(6) Friederich, E.; Kopiev, I.M.; Busalov, Y.E. and Weiss, G.Y. Fizica i Chimica Obrabotki Materialov, 6 (1975) 115-119.

(7) Madronero, A.; Prensa, M. and Sanzo J. Rev. Soldadura, 17 (1987) 115-128.

(8) Barnes, J.A.; Kumosa, M. and Hull, D. Composites Science and Technology, 28 (1987) 251-267.

Table: PARAMETER ADJUSTMENT OF TENSILE PROPERTIES AS A FUNCTION OF THE MANUFACTURING PROCESS (ideally .S = 1 P = 1)

Stainless steel Pickling Pickling Process for Sintering ~ilament without in in

any surface 33'\. acid 75'\. acid treatment

Deep cold compactation in 0,42 0,59 0,39 atmosphere, followed by a short hot pressing under vacuum p 1,22 1,39 1,18

Hot pressing under vacuum; S 0,59 0,69 0,46 constant pressure and moderate slope heating

V 0,76 0,80 1,18

22

2 I

20

19

18

O"'c (M Pal Ec:u (01.) tCE(O,o Ec (GPa) I '] I }-26~---------

•• e.e e .0 .eel eEM. P33.EE I .EM.P1S.EE •

BEM.P33.EB II··~ .EM.P1S.EB • x II ..... ,,,,

•• A,. .EM.P15.VM A,. 'EM.P33.VM • • • • A

• • • • • x ...... ~ A EM.P33.UT A

x x EM,P1S.UT A

A • A A A • x x •

2 3 SINTERIN6 TIME (hr,

rs 0 24 + EM.P1S.VM -EM.P33.VM

32 o o o

~ 21

&

"5 3

1'5 0

22 20

II 18 1&

14 26 a KEV: 12

a I iii - Infil~rated I a - lIIatrlx 2£

22

A IM.P33.VM a IM.P15.VM

A a A a A a

a • to _

• a EM: Elec:t~oplated

00 lIIatrlx ~C

lID. ..IM.P33.UT.. ..r • IM.P1S.UT _...... 0

t .... • ... "· : II 16

14

o P33= Picklin, In 33"10 acid o • EM.P33.UT

• o

a P7S:Picklin, in • EIot.Pl!i.UT 7S"oacid • "

V M:Vounl's lIIodulus : : : • UT :Ultimat. t.nsit.

• o

5 tre... . . L-,.L..----....L2-----'3L...J EE=Elonlatl~nat SINTERIN&TIME (h)

• la.1 tic hmll ECE r

.Ionlotlon Ecu (calculated filllr .. , 11.1114. VM

200 110

160

140

EB: Ultlmo~. Th.oreticalvalull {I.EM.VM

III.EIot.UT IV.IM.UT

FIG. 3.-MECHANICAL PROPIERTIES AS A FUNCTION OF MANUFACTURING PROCESS PARAMETERS

EV:.IIot.P7S.F AEM.P33.F ... 1M: Infiltrated matrix 320 cEM.P15.F AIM.P33.F

0 180 'E EM: EI.ctroplated matrix

300 ~ x EM.P15.1 -IM.P7S.1 a .... N .,.0 EM.Pll.1 01 M.P33.1 a .>< P33 = Pickl in! in 33.,. acid IE 280 ·0 0- 170 ;U

P7S=Pickling in7So'.ocid .... • 0 D • w .>< 260 ~x~o • 12: F = Flexure tlSt -u a .- :::l I: losiPlStu's tut <!) 240 x 0 a. 160 )(

x 0 ... a. -' 0- 220 x o ii- I!. 1/1 W

x D. 0 150 0-0- 200 x. 0 Z

III 180 ~~--. 0 :, ':-'x' "fJi IL u

160 140 1/1 ~,-W .~ " W

IL 140 jt'" . w 12:

1/1 , ' ' ... x ~ 120 ' x • 130 0-

100 1 2 3 SINTERIN& TIME (hrl

FIG.4._FRACTURE ENERGV FOR THREE POINT BENDING AND 10SIPESCU'S TEST AS A FUNCTION OF THE SINTERING TIME DURIN6 SAIotPLE MANUFACTURING, SHOWI N& THE DIFFERENCE FROM ELECTRODEPOSITION AND INFILTRATION AS A PROCEDURE TO APPORT THE ZINC MATRIX

197

198

hr. 2 hr. J hr.

Fig.1.- Microestructural changes in electrolytic linc matrix du­ring sintering

Fig.2.- Interface formation when the matrix was apported by elec­trolytic process.

soft steel hard steel Zn/SS (back face)

Fig. 5. - Feature of piercing of stainless steel filaments reinfor-', ced zinc plates in comparing with steel.

CREEP RUPTURE OF 1100 SERIES Al!SiC PARTICULATE MMC'S

s. PICKARD, B. DERBY

Department of Metallurgy & Science of Materials University of Oxford - Parks Road Oxford OX1 3PH OXFORD - England

ABSTRACT

The temperature dependence of creep in Al/20% SiC composite has been studied at 150°C to 350·C. The material shows a high stress sensitivity and large temperature dependence of creep strength. The stress levels required for creep rates of 10-~s-1 - 10-7s-1 lie above the first deviation from linearity of the tensile stress strain curve. High temperature creep rupture occurs by failure of the Al matrix. Voiding about the SiC reinforcement is seen at lower temperatures.

INTRODUCTION

Under creep testing conditions metal matrix composite materials containing discontinuous SiC, Whisker or platelet dispersions exhibit much higher creep resistance than that of the monolithic parent matrix alloy. This ability of MMC's to retain high temperature strength has recently been utilised commercially to increase elevated temperature performance in such high temperature applications as diesel engine pistons. However under increased loading these MMC materials have low creep rupture strains of 1-5% at temperatures of interest between 200-350°C /1/. This accompanies the well known low ductility of MMC on room temperature tensile testing /1/. The stress sensitivity of these materials during creep is large, indicating a difficulty in slip processes both at low temperature and at higher temperatures, leading to the low creep failure strains. When the stress sensitivity is represented in the form E z Can then n takes a value of 20 for these MMC's as compared to typically 3-4 for the matrix alloy, this lower level being a typical value for easy dislocation glide mechanisms and climb control of creep rate /2/. The high stress sensitivity and temperature independent ductility

200

cannot be reconciled with any simple creep model involving continuous dislocation mechanisms. It is possible that local yield events about the SiC dispersion during creep could play an important role in initiating creep failure at low strains and this has been investigated by the present study.

EXPERIMENTAL

The material chosen for the study is a 1100 series 'pure' Al reinforced by 20% of '1200 grade' SiC grit (3~m mean particle size) supplied by BP Advanced Materials. Creep specimens of 3mm cylindrical diameter and 6mm gauge length were machined from the as­received material. The small creep specimen dimensions were chosen so that as many test pieces as possible were obtained from the limited supply of starting material. Creep strains during deformation at 150, 200 and 350°C were measured by an external extensometer pair accurate to elongations of ±l~m. Elevated temperature tensile tests were also performed using a small furnace mounted on a screw driven Instron at a strain rate of approximately 2x10-3s-l. Specimen surfaces were polished prior to testing with 3~m to ~m diamond paste to allow optical and SEM examination of the bright and mirror-like surfaces. TEM observations were conducted on thinned specimens obtained by electropolishing of transverse sections obtained as close as possible to the fracture surface. Occasionally ion milling was used to make the specimens more transparent to the electron beam. TEM observations were performed in a Jeol 200CX microscope at 200Kv operating voltage.

MECHANICAL RESULTS

The creep behaviour of the composite at the temperatures 150, 200 and 350°C is shown in Figure 1. The material shows a high stress sensitivity and the value of n, where E = Con takes a value 19-20 indicating the difficulty in plastic flow processes in the alloy. At all the temperatures the creep curves show a three stage behaviour of initial primary creep, followed by steady state and tertiary behaviour at the end of the test. The typical behaviour of the composite is shown by the creep curve in Figure 2. The primary creep is characteristic in materials in which easy dislocation glide can occur and is to be expected in the pure AI. The creep rate was recorded down to low rates of below 10-7 with no indication of a threshold value which is common in similar materials with high stress sensitivity such as ODS materials /3/. The threshold limit is normally thought to correspond to the critical stress for a yielding event to occur in the matrix and below this threshold level a slow creep rate or no creep occurs because of the strengthening by the reinforcement. To determine the creep stresses at which yielding occurs in the composite the tensile stress strain curves at elevated temperature were studied at quasi-static strain rate 2xlO-3s-1 (Figure 3). In all cases the creep stresses used were above the yield point for the composite as determined by the initial deviation from linearity of the tensile curve. Lower stresses than these would need to be studied to determine if any true creep occurs below this threshold value.

201

It is noteworthy that even beyond the yield point the rate of creep is surprisingly slow indicating the plastic constraint due to the reinforcement. The tensile tests showed that the UTS of the composite decreased rapidly with temperature and this accounts for the rapid decrease in creep strength with temperature that is seen in the data.

The work hardening rate of the alloy on tensile testing shows the same trend as the UTS with an extremely high rate of hardening at room temperature and a reduction to a nearly negligible rate at 350°C prior to the onset of necking. The values of Youngs modulus could not be accurately obtained at the high temperature tensile tests because of the difficulty of obtaining elevated temperature strain measurements. However the rapid decrease in modulus recorded by the cross-head movement indicated a drastic loss in modulus at around 350°C so that the composite is no longer strongly influenced by the stiffness of the reinforcement phase and the matrix properties now determine the behaviour. The creep failure strain was about 14% at all temperatures and is slightly below the value recorded by the tensile test of about 20% at all temperatures. Some elevated temperature degradation of properties might have occurred. The high failure strain on room temperature tensile testing is surprisingly large and is not found in higher matrix strength composites which have been tested previously /2/.

METALLOGRAPHY

1.1. Fracture Surfaces

Failure of the composite invariably occurred by localised deformation in a necked region close to final fracture due to the high stress sensitivity of the material. Polished creep specimens showed intense slip during final fracture and fewer slip lines are seen further away from the necked fracture surface. SEM surface observations often showed the opening up of grain boundaries at high temperature consistent with the matrix rupture mechanisms (see Figure 4). The fracture surfaces after creep rupture were compared with the uniaxial tensile failure surface at room temperature so as to determine failure mechanisms. The room temperature test showed the expected ductile rupture with the fracture surface containing debonded silicon carbide particles siting within the ductile cusps as shown in Figure 5.' The rough surface of the particles suggested that some particle fracture might have occurred during failure. The creep failure at 350°C showed a different type of fracture surface with much less ductile voiding about the particles. Instead the fracture surface predominantly showed matrix ductile rupture between the particles with Al metal still adhered to most SiC particles siting on the fracture surface as shown in Figure 5. The fracture surface at 200°C was intermediate between the two already described. This failure behaviour is in agreement with the change in work hardening rate from high to low values at elevated temperature, so that the conditions for voiding about the particles can no longer be easily satisfied. At elevated temperature it appears that the matrix properties determine behaviour.

202

1.2. TEM observations

The untested composite showed a fine grain size of about 2~m. The SiC particles ranged in size and were often distributed inhomogeneously in the matrix mainly at grain boundaries. In the fine grains there were few stray dislocations and this is consistent with recovery and recrystallisation occurring during the high temperature processing stage of manufacture. Oxide films, apparently not broken down by the extrusion stage, were often observed about the SiC particles and a small number of inclusions were seen within the grains. High temperature creep specimens revealed a substructure little changed from that of the untested material even though considerable straining had occurred within a necked region from which the specimen was taken. There was some small tendency to form a cell substructure in some grains close to fracture consistent with the occurence of primary creep (See Figure 6). This poorly developed substructure would indicate that dynamic recovery has occurred at the high temperature and explains the lack of void formation. Examination close to the fracture surface after tensile failure occurred, showed large dislocation multiplication within the grains nearby. The presence of much smaller sub grains structure seen within the fine grains is probably explained by the formation of deformation subcells close to the fracture surface where high rates of strain hardening occur. Void formation was seen in the cross sections taken close to fracture consistent with the fracture surface observations.

DISCUSSION

The elevated temperature behaviour of the composite material is characteristic of particle strengthened materials in general where the reinforcement provides obstacles to the dislocations motion. The high stress sensitivity on creep appears to be caused by the material being tested beyond the yield point of the alloy so that accelerated yield and flow past the SiC particles occurs for small stress increases. The failure strain is determined by tensile rupture and no special creep failure mechanism occur. Easy matrix flow at elevated temperature mean that the failure of the matrix plays an important part in final fracture. Prediction of the creep strength of the composite appears to be simply governed by the tensile stress strain conditions of the test and elevated temperature degradation of creep properties has only a small effect. The loss of hardening at high temperature causes both a drastic loss in composite UTS and a low creep strength. This temperature sensitivity, recorded in the past as a high activation energy for creep to occur, results from some loss in the capability of the reinforcement to provide effective obstacles to dislocations and possibly results from the onset of climb around the particles /3/. It produces a reduction in composite modulus as the soft matrix now governs the mechanical behaviour.

203

'IJ"

T=MC T=201!C

(J= 34·5~a

1a-s 0

~ 0 Strain 1%14

-xl' 0

~

'K]?

~ 3J JDIMPaI 50 00 ~ 10 15 EJ

time21:ll's 1 JJ 35 4Il

Fig.l Stress dependence of creep. Fig.2 Typical creep curve at 350·C

11.;

150

125

10 15 20 2\ Strain ("/01

Fig.3 Tensile curves Fig.4 Specimen surface at 350°C.

Fig.5 Fracture surfaces, RT tensile failure (left), 350°C creep failure right.

204

Fig.6 Creep substructure after failure at 350·C (0 = 36MPA). Note oxide film around the SiC particle on left side of the picture.

REFERENCES

1. D. McDanels, Met. Trans., 16A (1985) 1105.

2. T.G.Neih, Met. Trans., 15A (1984) 139.

3. J.W.Martin, Micromechamisms in particle hardened alloys, Cambridge University Press (1980), p.165-175 .

ABSTRACT

MICROSTRUCTURAL STABILITY OF FIBROUS COMPOSITES BASED ON

MAGNESIUM-LITHIUM ALLOYS

M. WARWICK, R.T.w. CLYNE

University of Cambridge Department of Materials Science & Metallurgy

Pembroke Street - CB2 30Z CAMBRIDGE - England

A magnesium alloy containing about llwt.% lithium is shown to be attractive as a matrix for composite materials. The degradation of carbon, alumina and silicon carbide in this matrix is examined, and the effect of lithium on silicon carbide monofilament and multifilament is examined using a vapour phase impregnation technique. Only silicon carbide whiskers are found to be stable in this matrix, although a barrier layer of yttria is shown to protect silicon carbide monofilament from attack.

I - INTRODUCTION

Many light alloys are being explored as matrices for metal­based composite materials. For example, magnesium has been reinforced with boron [1], silicon carbide [2], graphite [3] and alumina [4-6]. Among other publications are a number dealing with aluminium-lithium matrices [7], and at least one in which pure lithium is considered [8]: this latter study explored the chemical compatibility between molten lithium and various candidate reinforcements, reporting a variety of chemical attack phenomena. However, there is a shortage of information on the use of magnesium-lithium alloys in metal matrix composites.

The Mg-Li phase diagram is unusual. No intermetallic compounds are formed and the lithium body-centred cubic structure exhibits solid solubility for magnesium up to about 70 at.% (90 wt.%). An alloy of Mg-ll wt.%Li displays a combi­nation of high melting point (about 593'C), low density (about 1.4 Mgm-3) and very high ductility. Detailed work on the fabrica­tion and properties of Mg-Li alloys was carried out by J.H Jackson et al in 1949 [9]. This study highlighted the difficulties of strengthening materials based on the bcc matrix by conventional methods. Attempts to develop precipitation-hardening systems were frustrated by overaging effects at room temperature. This is apparently a result of very high vacancy mobility, which also encourages dislocation climb and minimises work-hardening. The material exhibits pronounced creep at room temperature .

206

It might be expected from the above that Mg-Li alloys would be attractive candidates for reinforcement with ceramic fibres or particles. In practice they present certain difficulties, not least in terms of fabrication route development and other problems related to their high chemical reactivity. In the present paper an introduction is given to some aspects of microstructural stability of composites based on these alloys.

II - EXPERIMENTAL

2.1 Composite Fabrication

Full details of the preparation of the alloys and composite materials are given elsewhere [10]. All of the composites described here were manufactured by squeeze infiltration [11]. The entire operation is carried out in a vacuum chamber under controlled atmosphere. various reinforcements have been employed, including some based on alumina, carbon and silicon carbide. In the case of the silicon carbide whisker-reinforced material, attempts were made to produce unidirectionally-aligned material from the as-cast planar random composites by extrusion processing. Following experimental investigations of the role of extrusion conditions in controlling the degree of fibre damage in aluminium­based composites [12], a combination of low die angle, high extru­sion temperature and modest extrusion ratio was used to produce good alignment of whiskers with relatively high aspect ratios.

2.2 Microstructural Examinations

Specimens were examined by optical, SEM, and TEM techniques, before and after various thermal and mechanical treatments. Most of the problems of specimen preparation are those that arise from the chemical reac~ivity of the system. Exposed surfaces corrode quite rapidly, particularly in the presence of moisture. In particular, lithium loss from the matrix can be significant at room temperature. (For example, an initially single phase matrix can, when examined by x-ray diffraction - sampling the superficial 50~ or so - give rise to detectable hcp phase peaks on exposure to air for a few days). Preparation of TEM specimens is particularly hampered by such lithium loss: conventional ion beam thinning results in complete conversion of the thinned region of the matrix to the hcp phase. This effect can be reduced by carrying out the ion beam milling at low temperature, but the foil is then prone to oxidation during transfer to the microscope. Nevertheless, useful examination of any changes induced in the ceramic reinforcement is possible by TEM techniques. Deep etching to reveal fibre distributions by SEM examination was carried out by using 1% sulphuric acid as the etchant.

2.3 Lithium Vapour Impregnation

In order to isolate certain effects thought to be associated solely with the presence of lithium, a stainless steel vessel was constructed, designed to expose fibres to Li vapour at a selected temperature. The vessel, containing a support system for the fibres and some small pieces of Li, was evacuated and sealed before being placed in a furnace. Typically, a treatment period of 12 hours at 700'C (corresponding to a Li vapour pressure of lrnbar) was employed.

207

III - RESULTS AND DISCUSSIONS

3.1 Carbon Multifilament (in Planar Random Felt)

Little attention was devoted to carbon, as it soon became clear that extensive carbide formation takes place very rapidly on exposure to a Mg-Li melt. Fig.1 shows the microstructure in the as-fabricated form. The fibres, which were originally of circular cross-section, have been extensively attacked, and lithium­depleted (hcp phase) regions of matrix created around them. X-ray diffraction data indicated that Li2C2 had been formed. A degree of fibre cracking was also observed.

3.2 Alumina Staple Fibre ("Saffil")

This fibre contains about 5 wt% Si02 and is primarily in the form of the ~-A1203 phase [13,14]. In the as-fabricated form there was little obvious evidence of gross chemical attack. There were, however, indications that the (silica-based) binder was being rapidly dissolved on contact with the melt and then carried through the preform, becoming concentrated in the lower regions. In these areas, certain precipitates, notably Mg2Si, often appeared in some profusion. In addition to the binder reaction effects, there were indications that a degree of interaction was taking place with the fibre. The silica in "Saffil" is concentrated on the free surface and at grain boundaries [14]. Reduction of the surface layer was probably responsible in part for the observed high fibre-matrix bond strength, but penetration of lithium down the grain boundaries also appeared to take place, leading to severe embrittlement of the fibres. For example, in Fig.2 it can be seen that stressing of these composites led to much more severe fibre cracking than is observed with other matrices.

3.3 Silicon Carbide Fibres

3.3.1 "Nicalon" Multifilament

This fibre is not pure SiC; the product used here has a very fine grain size (about 2nm), with about 15 wt% free carbon and 25 wt% free silica present in amorphous form at the grain boundaries [15]. Composites produced by axial infiltration of fibre bundles exhibited good matrix penetration, even into regions of very high fibre volume fraction, an example of which is shown in Fig.3(a). However, it became clear that this system was highly unstable microstructurally. In particular, the fibres tend to absorb lithium very rapidly. In Fig.3(b), which is a backscattered SEM image, the lithium-depleted regions appear light. It can be seen that these occur around the fibres: in areas of high fibre population the depletion has been sufficient to cause formation of the hcp/bcc eutectic structure.

This lithium penetration, which occurs during and immediately after infiltration, is sufficient to cause differential effects along the infiltration axis. Furthermore, although the change in matrix composition is exaggerated by the high volume fraction of fibre in the uniaxial bundles, the consequences for the fibre itself are greater at lower fibre loadings. For example, Fig.4, which shows structures at different heights in a planar random Nicalon preform, illustrates that the fibre degradation is pronounced. (This is less obvious near the base, where the melt contact time and high temperature exposure period are shorter.) The experiments in which the Nicalon fibre was exposed to lithium

208

vapour confirmed that the element is rapidly absorbed up to very high levels (~30 wt%), causing dramatic embrittlement.

3.3.2 "Sigma" CVD Monofilament (Tungsten Core)

In the case of.the large diameter CVD monofilaments, which are essentially pure SiC, there was no obvious evidence of pro­gressive chemical reaction taking place - for example new x-ray peaks did not appear. Nevertheless, the fibres did tend to suffer pronounced degradation - apparently in the fo~ of grain boundary attack. The grain structure in these fibres is of a relatively coarse, radial columnar type [20]. For example, in Fig.S a compar1son is presented between the fracture surfaces of individual monofilaments broken manually before and after exposure of the fibres to lithium vapour. In the latter case (when the fibre is significantly embrittled) there is a clear tendency for cracks to follow intergranular paths across the section and along the fibre length (trends also observed in the composite). It may be noted that monofilaments of this type can be protected against attack. For example, sputter-deposited diffusion barrier layers [17,18] can be employed to isolate fibre and matrix. That these can be effective in the present case is illustrated in Fig.6, which demonstrates that a Y203 layer of about 200nm thickness can provide good protection against the ingress of lithium.

3.3.3 "Tokawhisker" Whiskers

It was found that the (single crystal) whiskers did not appear to suffer any chemical reaction on contact with the Mg-Li matrix, even after prolonged exposure to high temperature. For example, the TEM micrograph shown in Fig.7 is of a specimen heated for about 12 hours at SOO·C. Although the matrix is fully hcp, having suffered extensive lithium loss during foil preparation, it is clear from Fig. 7 that the whisker surface has not been noticeably affected by this very severe heat treatment.

This observation should be considered in the light of thermodynamic information about the system [19,20]. Among possible chemical reactions, the two listed below would appear the most likely. (The free energy values are not strongly dependent on the temperature chosen.)

2Mg + SiC ..... Mg2Si + C (1)

AG~OOK =-7 (±lS) Jmole-1

2Mg + Li + SiC ..... Mg2Si + tLi2C2 (2)

AG~OOK --18 (±2S) Jmole-1

209

These data are consistent with the current observations, in which the apparent absence of even a very thin reaction layer suggests that it is the lack of a thermodynamic driving force, rather than slow reaction kinetics, which is responsible for the stability of the system. It may be noted that the grain boundary penetration observed with the "Sigma" fibres may have been encouraged by chemical effects, such as an affinity for traces of oxygen in the boundaries.

(3)

AG~OOK =-123 (is) kJmo1e-1

IV - SUMMARY AND CONCLUSIONS

The Mg-Li system exhibits a number of interesting features, in addition to its attractive combination of relatively high modulus and very low density. Many of these features are associated with the very high atomic mobility of lithium, even at relatively low temperatures. Among other effects, this leads to rapid coarsening of precipitates: in combination with rapid dislocation climb, this destroys the efficiency of conventional precipitation hardening mechanisms. The mobility of the lithium atoms, together with their high chemical activity, also affects the viability of mechanically incorporated inclusions such as ceramic fibres. Although a satisfactory fabrication route has been developed for making composites based on the Mg-Li matrix, the choice of reinforcement is severely restricted by the reactivity aspects. Basically, of the reinforcements commonly available in fibrous form, only silicon carbide whiskers remain stable: the others all undergo highly deleterious chemical attack and/or grain boundary penetration. Although the magnesium is often as active in chemical terms as the lithium, it is, in many cases, the very fast kinetics associated with the latter that is responsible for an unacceptable degree of degradation occurring during fabrication. It may be that certain mechanical performance requirements would be best satisfied with reinforcements that are continuous and of large diameter. For these cases, it appears likely that the development of barrier layer coatings on fibres may be desirable. Results presented here indicate that sputtered yttria layers are promising in this respect.

ACKNOWLEDGEMENTS

The authors are grateful to BP plc for funding the project within which the work described here has been supported. Particular acknowledgement should be made to Dr. C. Brown, Dr. A. Begg and Mr. J. Robertson, of the BP Research Centre, Sunbury-on-Thames, for active support and stimulating discussions.

REFERENCES

[1] T.W. Chou, A. Kelly and A. Okura, Composites, II (1985), 187-206

[2] B.A. Mickucki, 5.0. Shook, W.E. Mercer and W.G. Green, Conf. Int. Magnesium Association, Los Angeles, (1986)

210

[3] B.J. MacLean and M.S. Misra in "Mechanical Behaviour of Metal-Matrix Composites" (Proc. Symp.1982, J.E. Hack and M.F. Amateau, eds.), 195-212, Metallurgical Soc. of AlME, Pennsylvania, (1983)

[4] J.E. Hack, R.A. Page and G.R. Leverant, Metall. Trans A, l5A (1984), 1389-1396

[5] R.A. Page, J .E. Hack, R. Sherman and G.R. Leverant, Metallo Trans A, l5A (1984), 1397-1405

[6] J.T. Evans, Acta Met, .li (1986), 2075-2083 [7] D. Webster, Metallo Trans. A, .l3A (1982), 1511-1519 [8] R.T. Swann and D.M. Easterling, Composites, II (1984),

305-309 [9] J.H. Jackson, P.D. Frost, A.C. Loonam, L.W. Eastwood and

C.H. Lorig, Trans. TMS-AlME, ~ (1949), 149-168 [10] J.F. Mason, C:M. Warwick, J.A. Charles and T.W. Clyne,

"Magnesium-lithium alloys in metal matrix composites-A preliminary report", submitted to J. Mat. Sci, 1988

[11] T.W. Clyne and J.F. Mason, Metallo Trans. A, l.6.A (1987), 1519-1530

[12] C.A. Stanford-Beale and T.W. Clyne, "Extrusion and high temperature deformation of fibre reinforced aluminium", published in Comp. Sci. and Techn. (special issue on metal matrix composites), (1988)

[13] J.D. Birchall, Trans. J. Br. Ceram. Soc., ~ (1983), 143-145 [14] G.R. Cappleman, J.F. Watts and T.W. Clyne, J. Mat. Sci., 2Q

(1985), 2159-2168 [15] G. Simon and A.R. Bunsell, J. Mat. Sci., ~ (1984),

3649-3657 [16] R. Warren and C.H. Andersson, Composites, II (1984), 16-24

and 101-111 [17] R.C. Mehan, M.R. Jackson and M.D. McConnell, J. Mat. Sci.,

~ (1983), 3195-3205 [18] R.R. Kieschke, R.E. Somekh and T.W. Clyne, "Protection of

SiC monofilaments against attack in metal matrix composites by sputter-deposited barrier layers", in preparation

[19] "Janaf Thermochemical Tables, 2nd edition", US Department of Commerce, NSRDS-NBS, 37 (1971)

[20] C.E. wicks and F.E. Block, "Thermodynamic properties of 65 elements-Their oxides, halides, carbides and nitrides", U.S. Govt. Print. Off., Washington, (1963)

Fig. 1 Optical micrograph of as-fabricated Mg-12 wt. %Li/15 vol. %C planar random composite

Fig. 2

Fig. 3

Fig. 4

211

(a) (b) (c)

SEM micrographs of polished longitudinal sections of 12 vol. % "Saffil" composites after tensile testing (to a few per cent strain) (a) & (b) with a Mg-12 wt. %Li matrix and (c) with an Al-2·S wt.%Mg matrix

(a) (b)

Transverse sections of uniaxial Mg-12 wt. %Li/ 60 vol. % "Nicalon" composite: (a) Optical micrograph of high fibre content area and (b) backscattered SEM micrograph of a region of variable fibre population

Optical micrographs of Mg-12 wt.%Li/20 vol.% "Nicalon" planar random composite parallel to the infiltration axis (a) about 3mm above the base and (b) about O·Smm above the base

212

Fig. 5

I!!>" ' ..... 'I .t, ~,~:f ;, ' J.

" ...... 1 ...... ) '\ I ,

'. ~ '\ .' , ... r '" \',. " , I

Fig. 6

Fig. 7

(a) (b)

SEM micrographs of the fracture surfaces of individual "Sigma" fibres broken by hand (a) in the as-received condition and (b) after impregnation with lithium by exposure to lithium vapour (about 1mbar for 18 hours at 700'C)

l -• . , , " \ I

, '. L . .

(a) (b)

Back-scattered SEM images of single "Sigma" fibres embedded in a Mg-12 wt. %Li matrix, after heat treatment for 1 hour at 300'C, with the fibre (a) in the as-received condition and (b) pre-treated by sputter­deposition of a layer of yttria

(a) (b)

TEM bright field micrographs of whiskers in a Mg-,12 wt. %Li/20 vol. % "Tokawhisker" planar random composite (a) as-fabricated and (b) after heat treatment for 280 hours at 500'C

MICROSTRUCTURAL DEVELOPMENT AND MECHANICAL BEHAVIOUR OF SiC

WHISKER-REINFORCED Mg-Li ALLOYS

J. MASON, R.TW CLYNE

University of Cambridge Department of Materials Science & Metallurgy

Pembroke Street - CB2 3QZ CAMBRIDGE - England

ABSTRACT

Preform production, squeeze infiltration and extrusion procedures are described for the preparation of single phase (~) Mg-Li alloys reinforced with SiC whiskers. Observations are reported on the generation of whisker alignment and second phase distribution as a result of preform binder dissolution and re­precipitation.Tensile test data are given, showing that the presence of whiskers can lead to considerable property enhancement. Dynamic stress relaxation processes appear to affect the tensile behaviour at room temperature over the range of strain rates employed during conventional testing. It is thought that this is associated with the very fast diffusion kinetics exhibited by the lithium.

l. INTRODUCTION

Composites based on magnesium, offering potential for attractive combinations of low density and high stiffness/strength, have been under study for some time [1]-[3]. Among magnesium alloys now attracting particular interest in this context are those in which significant levels of lithium are present, which can exhibit extremely low densities. Such alloys flrst received study some years ago [4], but only recently have some of the problems of introducing flbrous reinforcement been explored [5],[6]. It has been found [5] that, although many flbrous reinforcements undergo degradation in contact with Mg-Li alloys at modestly elevated temperature, SiC whiskers tend to remain unaffected. In the present work, some observations are presented on the development of certain microstructural features during preparation of SiC whisker-reinforced Mg-Li alloys. A brief examination is then made of the behaviour of these composites during tensile loading and some simple points are identifled about optimization of performance.

2. EXPERIMENTAL

2.1 Composite Fabrication

The whiskers employed were the "Tokawhisker" product manufactured by Tokai Carbon (Japan). In the present work, the alloy composition was about 11 wt%Li, which is in the fully ~ (bec) regime. Composite castings were produced by

214

melt infiltration of planar random whisker preforms under controlled atmosphere. These castings were then machined into cylindrical billets (with the cylinder axis lying in the plane of alignment) and extruded under selected conditions, also under controlled atmosphere. The whiskers were fIrst dispersed in water containing a colloidal silica-based binder. The slurry was then compressed and dried to give an approximately planar random array of whiskers, bonded at contact points. In the present work, conditions were chosen so that the whiskers filled space in the preform to about 20 % by volume. In addition to preforms produced in this way, commercial products supplied by Hepworths Ltd were also employed. The preform, of 75 mm diameter and 40 mm thickness, was placed in a die heated to 420 C and liquid metal at 900 C was poured onto the top surface. Pressure was applied via a hydraulic ram set to a maximum pressure of 30 MPa. The onset of inflltration was monitored by means of strain sensors on the rods supporting the die assembly and taking the thrust of the applied load: melt entry is detectable as a distinct "blip" on the rising curve representing the melt pressure history. Entry occurred at a pressure of about 5-6 MPa in the present work. This may be compared with theoretical predictions [7], based on an idealized fIbre array, of about 2-4 MPa, obtained using values thought to be broadly appropriate for the melt surface tension and parameters describing the fIbre assembly geometry. It was found that the preforms suffered a permanent compression of about 10 % during inflltration. In an attempt to minimize whisker fracture during extrusion [8], the process was carried out at a high temperature of about 450 C, using a die with a low semi-angle of 20 degrees and an extrusion ratio of 10.

2.2 Metallographic Techniques

Specimens were examined by optical, SEM, and TEM techniques, before and after various thermal and mechanical treatments. Most of the problems of specimen preparation are those that arise from the chemical reactivity of the system. Exposed surfaces corrode quite rapidly, particularly in the presence of moisture. In particular, lithium loss from the bec matrix can be signifIcant at room temperature. (For example, an initially single phase matrix can, when examined by x-ray diffraction - sampling the superficial 50 ~m or so - give rise to detectable hcp phase peaks on exposure to air for a few days). Preparation of TEM specimens is particularly hampered by such lithium loss: conventional ion beam thinning results in complete conversion of the thinned region of the matrix to the hcp phase. This effect can be reduced by carrying out the ion beam milling at low temperature, but the foil is then prone to oxidation during transfer to the microscope. Nevertheless, useful examination of any changes induced in the ceramic reinforcement is possible by TEM techniques. Deep etching to reveal fIbre distributions by SEM examination was carried out by using 1 % sulphuric acid as the etchant

2.3 Mechanical Testing

Tensile testing was carried out with a screw-driven machine, using standard HoundsfIeld test piece geometry. Strain rate control was effected via the crosshead speed. Strain was monitored by means of strain gauges attached to the gauge section by adhesive. Tests were carried out over a range of temperature and strain rate, using both as-cast and as-extruded material. For the former type of specimen, the tensile axis lay in the plane of whisker alignment. while in the latter case it was parallel to the extrusion axis. The test data given refer to composites from preforms produced within the programme using optimised dispersion techniques. (Material from preforms containing poorly-dispersed whiskers, and from the commercial preforms, gave inferior mechanical properties.)

215

3. RESULTS

3.1 Microstructural Features

It was found that homogeneous distributions of well-aligned whiskers could be produced under suitable conditions. However, the initial dispersion step can be critical in this regard, in that whisker agglomerates from the as-received material which survive this stage tend to persist through the casting and extrusion operations. The microstructures shown in Fig.t resulted from procedures identical except for the degree of agitation of the whisker slurry. It is clear that plastic deformation, including the rotational component of the plastic strain field responsible for fibre alignment, is not impressed onto dense whisker agglomerates during the extrusion, so that inhomogeneities and regions of poor alignment were produced.

It was noted that a quantity of particulate matter was present in addition to the aligned whiskers. Although some of these particles could be traced to the as­received whiskers, a number of relatively large (1-3J.lI11) spherical precipitates were present which were of different origin. It was found that these had a distinct "core and sheath" structure, evident in Fig.2. The cores were in many cases found by energy-dispersive X-ray analysis to correspond to CoSi, or in a few cases to silicides of other transition metals such as Fe. These cores were found to correspond to particles present in the as-received whiskers, believed to have been produced during the process of whisker growth [9]. The outer sheath was in all cases found to correspond approximately in composition to Mg2Si. It seems probable that these coatings resulted from initial dissolution of the silica binder in the alloy melt, followed by precipitation of this silicide on the core particles during solidification. Fmally, it may be noted that, although the whiskers showed no sign of interfacial chemical attack by the matrix (even during prolonged heat treatment [5]), there was, in the case of the composites produced from commercial preforms, some evidence of fine scale deposition of precipitates on the whisker surfaces. Fig.3 is a TEM micrograph of as-cast composite, showing decoration of whisker facets with particles of about 20 nm diameter. Although the chemistry and structure of these have not yet been established, the dark field image shown indicates sets of them on two of the whisker facets to have a common orientation, suggesting either epitaxial growth on the whisker surface or possibly shape anisotropy and a preferred collision and adhesion geometry. These precipitates presumably also result from binder dissolution and reprecipitation effects. They were not observed with the preforms produced within the cqrrcnt programme.

3.2 Mechanical Behaviour

Some typical stress-strain curves from the tensile tests are shown in Fig.4, which gives a very brief indication of the effects of (a)fibre alignment, (b) strain rate and (c) test temperature. A guide to the nature of the failure in room temperature tests is given by the two fracture surfaces shown in Fig.5, which are from specimens tested at different strain rates. These both reveal local dimples around exposed whisker ends and limited pull-out, although these are less pronounced for the higher strain rate specimen, which also shows a few examples of large precipitates on the fracture surface.

216

4. DISSCUSSION AND CONCLUSIONS.

The results presented suggest that improved tensile strength and ductility are obtained with a homogeneous distribution of whiskers exhibiting high aspect ratio and a degree of alignment along the tensile axis. Suitable processing conditions have been outlined, allowing extrusion with relatively little whisker fracture. Although prediction of the deformation conditions under which discontinuous fibres will fracture is a complex matter, it is possible to attempt semi-quantitative modelling. For example, a model based on the concept of a critical strain rate above which dislocation climb will be insufficiently rapid to prevent pile-ups at fibres [10],[11] has been used to predict the onset of severe damage [11], [12]. The Mg­Li alIloy matrix used here is known to exhibit very rapid Li diffusion, although diffusivity data in the literature appear to be limited to a single value of 10-10 m2 s­I, for a temperature of 420 C [13] - which happens to be close to the extrusion temperature employed in the present work. Although it is unclear which diffusivity value is appropriate to employ in the climb model for the case of an alloy, substitution of this value in the expression for the critical strain rate to cause fibre fracture [11],[12] leads for these whiskers to a very high value of about 103 s-I, assuming volume diffusion to be dominant. Estimation of local strain rates during extrusion is very difficult, although measurements made with other MMC systems [8] suggest that, for the die shapes and conditions employed here, the peak value is almost certainly not more than about 10 s-l. These estimates are consistent with the observation of little whisker fracture: in fact, the very rapid diffusion in the alloy at these temperatures may be responsible for the failure to redistribute whiskers located within agglomerates, as it allowed matrix stresses to remain very low. At room temperature, the critical strain rate regime will be much slower, and might perhaps be deduced as lying between the two values for the tensile test curves in Fig.4 (a).

With regard to other microstructural features, it is clear that good alignment of whiskers along the tensile axis is beneficial, as is minimization of the incidence of extraneous precipitates. (All attempts to induce strengthening by age-hardening mechanisms were frustrated by over-ageing at room temperature [4].) It would appear that the stress relaxation kinetics are important in determining the nature of the fracture process. However, unequivocal identification of the dominant relaxation mechanisms under different regimes of temperature and strain rate requires a full programme of steady state creep characterisation, which is currently under way.

Acknowledgements

The authors are grateful to BP pic for funding the project within which the work described here has been supported. Particular acknowledgement should be made to Dr. C. Brown and Mr. ] . Robertson, of the BP Research Centre, Sunbury-on-Thames, for active support and stimulating discussions. Mr.I.A.G.Fumess, of Cambridge University, was involved in some of the work on precipitate identification.

217

References

[1] BJMacLean & MS Misra, "Thermal mechanical behaviour of graphite/magnesium composites", 195-212, Proc Symp. "Mechanical Behaviour of Metal Matrix Composites" (1982), JF Hack & MF Amateau (eds), Met Soc AIME, Warrendale Pa (1983)

[2] JE Hack, RA Page & OR Leverant, "Tensile and fatigue behaviour of aluminium oxide fibre reinforced magnesium composites: Part I-Fibre fraction and orientation", Metall. Trans A,llA (1984), 1389-1396

[3] IT Evans, "Fracture and subcritical crack growth in alumina fibre/magnesium composites", Ada Met, ~ (1986),2075-2083

[4] JHJackson, PDFrost, ACLoonam, LWEastwood & CHLorig, "Magnesium-lithium base alloys-Preparation, fabrication and general characteristics", Trans. TMS-AIME, ill (1949), 149-168

[5] JFMason, CMWarwick, JACharles & TWClyne, "Magnesium-lithium alloys in metal matrix composites-A preliminary report", submined to J. Mat. Sci, 1988

[6] CMWarwick & TW Clyne, "Microstructural stability of fibrous composites based on magnesium-lithium alloys", to be published in ECCM3 (1989)

[7] 1W Oyne & JF Mason, "The squeeze infiltration process for fabrication of metal matrix composites", Metall. Trans. A, 18A (1987), 1519-1530

[8] CAStanford-Beale & TWClyne, "Extrusion and high temperature deformation of fibre-reinforCed aluminium", to be published in Compo Sci. and Techn. (special issue on metal matrix composites), (1988)

[9] A. Yamamoto/Tukia Carbon. 'Process for Preparing Silicon Carbide Whiskers.' Patent GB21lS33A (1983).

[10] FJHumphreys & PNKalu "Dislocation-particle interactions during high temperature deformation of two-phase aluminium alloys", Acta Met II (1987) 2815-2829

[11] FJHumphreys "Deformation and annealing mechanisms in discontinuously reinforced metal matrix composites", in Proc 9th Riso Int Symp Met & Mat. Sci., SlAnderen et al (eds), p.51-74, Riso Nat. Lab., Denmark (1988)

[12] CAStanford-Beale & TWClyne "Deformation of Fibrous metal matrix composites at temperatures close to the matrix solidus", in Proc 9th Riso Int Symp Met & Mat. Sci., SIAnderen et al (eds), p.479-484, Riso Nat. Lab., Denmark (1988)

[13] Y. Iwadate, M. Lassourani, F. Lantelme and M. Chemla Electrochemical Study of Mass Transfer in Li - Mg and Li - Mg - Al Alloys. J. App. Electrochem. 17, (1987), 385 - 397.

218

Fig. 1

Fig. 2

(a) (b)

SEM micrographs of extruded composites. The surfaces are deep etched in a plane parallel to the extrusion direction and show: (a) a dense ball resulting from inadequate whisker dispersal in the original preform and (b) a homogeneous and well alligned whisker distribution.

A backscattered electron image showing the phase structure of a binder derived precipitate.

StressMPa

(a>

(b) Stress MPa

(c> Stress MPa

Fig. 4

400 .,---------------.,

Extruded 300

200

100 Unreinforced

O~--~--~--~~--~------~ o 2 3

StraiD CJI,

400,-------~--------------~

300

200

100

O~------~------_r--~--~ o 2 3

SllIIn ...

400,-------------------------~

lOOK 300 205K

200

100

o 2 .3

SIraln%

Stress - strain curves. derived from tension tests showing the effects of (a) whisker allignment, (b) the strain rate and (c) the temperature, upon the work hardening characteristics of the Mg - 12 wt% Li +20 vol% SiCw composite.

219

220

Fig. 3

Fig. 5

Bright and dark field transmission electron micrographs showing a fine dispersion of particles at the whisker -matrix interface.

(a) (b)

Representative fracture surfaces from tensile tests at ambient temperature and (a) a strain rate of 104 s-l and (b) 10-2 sol.

Abstract

PARTICLE REINFORCED MAGNESIUM ALLOYS

J. SCHRODER, K.U. KAINER, B.L. MORDIKE

Institut fUr Werkstoffkunde Technical University of Clausthal Agricolastrasse 6, 3392 CLAUSTHAL-ZELLERFELD - West Germany

Particle reinfored MMC's are interesting in application in arduous enviroments. Advantages of these composites are increased modulus, strength, high temperature properties and wear restistance. The thermal expansion is reduced. In this paper the microstructures and properties of composites with different particle additions, eg SiC, TiB2 , Ti(C,N), AIN and Al203-platelets produced by powder metallurgy techniques are dicussed.

I. Introduction

Magnesium base materials are ga~n~ng in importance due to their low density for aerospace, space, military and automobile appli­cations. The ever increasing demands on these materials in recent years led to the development of high performance materials. Conventional material developments using precipitation and solid solution hardening and grain refinement was unable to eliminate some of the disadvantages of magnesium alloys as for example the low modulus, poor wear resistance, poor high temperature strength and high coefficient of thermal expansion. Particle and fibre reinforced magnesium matrix composites can, by suitable selection of matrix and additives, exhibit a combination of metallic and ceramic properties. such a property profile opens the door to new applications such as bearing materials, pistons, gudgeon pins etc. The following methods can be used to manufacture particle rein­forced magnesium materials: casting by stirring in the particles (P. K. Rohatagi et al./I/) or by powder metallurgical techniques

222

(R. L. Trumper /2/, L. Erich /3/). Problems in the first method arise due to the reaction between the fibres and the matrix melt and also in the dispersion of the particles. In the powder metal­lurgical techniques eg consolidation by extrusion, it is possible to largely avoid chemical reactions. There are, however, greater problems in obtaining a satisfactory dispersion of the particles. A homogeneous dispersion is only possible if special dispersion methods are employed. The preparation and properties of particle reinforced magnesium materials are discussed below. The develop­ment of microstructure and the properties of a AZ91 with SiC, TiB2' AIN, Ti (C, N)50:50 and Al 0 -platelets are presented and compared. The matrix was prepare~ ~n form of powder by machining and attrition.

II. Production of Composites

The AZ91 powder was prepared by Mimeta S. A. Lausanne using mecha­nical methods. The fine powder with a maximum particle size of 63 ~m was seperated out and mixed with the additive. The additive was either SiC-, AIN-particles or A1 20 i -platelets with a mean size of 6 ~. Table 1 gives the chemical composition of the powder.

____ ~Al -"-_--'Z::.;n.:........,,.....-__ Mn=-=r_-::-"-S::-i -;;r_-::-"C-::u,--;;;_-: Be 9,5 % 1,0 % 0,30 % 0,3 % 0,2 % 15 ppm ma.x.

Fig. 2 shows the size distribution of the powder used. The physical properties of the additives and the dSO values are collated in Table 2. Dispersion was carried out dfy in a mixer followed by mechanical agglomeration in a ball mill.

particle SiC TiB2 AIN Ti(C,N) A1203-Platelets crystal type hex hex hex hex diameter (~) 6.5 5-7 5.0 6 6.2 density (g/cm 3 ) 3.21 4.51 3.26 5.18 3.90 Mohs hardness (max.13) 9.7 9.5 7.0 9.0 coeffcient of_5h~~al 4.7 4.6- 5.5 8.4 4.6 expansion (10 K ) 6.4

Table 2: Properties of particles added.

The flow dia.gram, Fig. 5, shows the production route of particle reinforced composites. The compact materials were produced by powder metallurgical techniques. Extrusion was used to consolidate the powder.

III. Microstructures and Properties of Composites

The microstructures of the materials after consolidation are shown in Fig. 4. With platelets, alignment in the direction of extrusion is possible (J. A. Black /4/). On the other hand, particle rein­forced material showed no difference between and the longitudinal and transverse directions. The mechanical properties of the particle reinforced composites with magnesium matrix are shown in

223

Fig. 6. With an addition of 15 vOl.-% of different particle reinforcements increasing strength and hardness was observed. Only the composites with A1N-particles showed a decreased strength. Wetting problems, chemical reactions and the powder size distri­bution are reasons for this behaviour. The youngs's modulus, Fig. 7, and wear resistance, Fig. 9, are increased with the particle addition. An decreasing thermal expansions coefficient, Fig. 8, was observed.

IV. Conclusion

The results show that is possible to produce particle reinforced magnesium using P/M-techniques with interesting properties. Particle reinforced composites represent one of the most inexpen­sive promising materials for automotive applications and have been sucessfully tried out as bearings, pistons and cylinder liners. Composites offer an improvement of the mechanical properties and wear restistance, but there are many problems during product.ion and further working on these materials eg the reactivity of magnesium, the dispersion of particles depending on the process conditions, and wetting problems. It can be seen that not every addition material is suitable for particle reinforcement . Never­theless homogenous distribution and increasing mechanical proper­ties, wear resistance and reduction of the thermal expansion coefficient are possible, using appropriate P/M techniques.

V. References

/1/ P. K. Rohatagi, R. Asthana, S. Dee; Solidifiation, structures and properties of cast metal-ceramic particle composites, Int. Metals Review (1986) 31, 3, 115 - 136.

/2/ R. L. Trumper; Metal MatriX-Composites - Applications and Prospects, Metals and Materials (1987), 662 - 667.

/3/ L. Erich; Metal-Matrix Composites, The Int. J. of Powder Metallurgy (1987) 23, 1, 45 - 54.

/4/ J. A. Black; Shaping Reinforcements for Composites, Advanced Mat. & Processes, Met. Process. 3 (1988) 51.

Fig. 1: SEM of magnesium powder AZ91.

25 ~tQ AZ91

Fig. 2: Size distribution of matrix alloy AZ 91.

224

Fig. 3: Particle morphologie of additives: a. SiC d. A1N b. TiE e. A1203-platelets c. Ti(6,N)50:50

Fig. 4: Microstructure of composites after extrusion. a. AZ91 + 15 vol.-% SiC b. AZ9l + 15 vol.-% TiE c. AZ91 + 15 vol.-% Ti(j,C)50:50 d. AZ9l + 15 vol.-% A1N e. AZ91 + 15 vol.-% A1203-Platelets

Iogllt ... tal aUo,s A1-.I1g-ba ..

mixed powdtrs

T

i l50

E. r ~

i i "5

aUoytd powdtrs

I

pttSsing cotd/hot

--r I

roU""l

I plat./pro'.

4191 4191 ·1SVol,·/.SiC

225

partidls p1alll.ts

... lIfIlI-tion

Fig. 5 : Flow chart mixing production process. T

hotUtatic preUIg

~ txlnJsion Iarging

~ T profil cllllpOntnl

• V'Kklts hardnt ..

o ttnsil. stttrqth

o "ongat..,

4191 AZ91 41" AZ" ·lSVal.%AIZ03 -pla'''tts

'lSVol.% Ti82 'ISVol·I.Ti!C.NI -lSYot.%AIN

of

Fig. 6: Mechanical properties.

226

., .. .£

I

I Young's modulus I

I thermal e.pansion coefficient I

AI" AI" AI" AI" Al" AI" -'6\1at-t.s.c:: ·'IS'Yd·I.l~ ·~VooI9f.TtC)I.'&1id,%.lIN ·1'5VIlll·/ ..... ~ ­OIIot4'leh

I!t2Ua. paral!!t'tr 5 RPI'1 (pin on di$c) : 250 min-1 load 7 N pin diameter 6mm retarding roller 42(rMoV no lubrication

Fig. 7: E-modulus of particle reinforced materials (various additions).

Fig. 8: Coefficient of thermal expansion.

:i.20 .. ~

ii :.

10

o AZ 91 (ex trusion) 83 HV 10 [] AZ91+Si( 13SHVlO x AZ91+TiB2 14OHVlO + AZ91+Ti(C,N) 150 HVlO • AZ91+AIN 130HVlO • AZ91.AI2~ - platelets

140 HV 10

°0~--------~20~-------4~0~------~~---------~80~----~

timelmin ) -

Fig. 9: Hardness and wear.

HEAT-TREATMENT EFFECTS IN S -ALUMINA FIBRE REINFORCED ALUMINIUM ALLOY 6061

ABSTRACT

C. FRIEND, R. YOUNG, I. HORSFALL

Royal Military college of Science (Cranfield) Shrivenham - SN6 BLA SWINDON - England

This paper presents the results from an experimental investigation into the age-hardening response of short d alumina fibre reinforced 6061. It is shown that chemical effects degrade the age-hardening response of these composites. These effects result from subtle interactions between the active matrix alloying elements and the surfaces of the alumina fibres. Such interactions degrade the age-hardening potential of the matrix and reduce the peak hardness attainable from the composites. It is also shown that chemical interactions between the preform binder and matrix alloy can result in anomalous age-hardening

INTRODUCTION

The potential of Metal Matrix Composites (MMC) have been recognised for many years since alloys reinforced with strong stiff fibres represent a unique method for tailoring mechanical properties to particular applications. However most evaluations of these materials have concentrated on optimising the properties by changing the type of fibre, its size, and orientation. Whilst these factors are important, it has recently been shown/ 1,2/ that in short-fibre reinforced MMC the properties of the matrix can also contribute to the final composite properties. The matrices of MMC can therefore constitute an 'active' component which can be used to alter the final properties of the composite by heat-treatment.

Most work on MMC has concentrated on light-alloy systems where the most common 'active' matrices are age-hardening alloys. However when heat-treating age-hardenable MMC it is often assumed that the matrix responds in a manner identical to the unreinforced alloy. Little consideration is therefore given to the presence of the fibre array and its effect on the microstructure and properties of the heat-treated matrix. There is little

228

published data describing the age-hardening response of alloys containing fibre arrays, however the evidence that does exist suggests that the in-situ age­hardening of the matrix can be considerably altered by the presence of reinforcing fibres/3,4,5f, These effects are believed to result from physical interaction phenomena and are similar to the processes observed in SAP/6,7/. In liquid metal processed MMC the fibre array can also potentially alter age­hardening through additional chemical phenomena. These can arise from alloy segregation effects/Sf or from matrix fibre reactions. These result in either the formation of unexpected second phases/9/ or the depletion of alloying elements from the matrix/lO/. The nature and extent of such effects depend on the actual composite system but must always impair hardening if the active age-hardening elements are involved in the interactions.

Recent work/5/ has identified the nature of the physical interaction phenomena in short o-alumina fibre reinforced 6061. This paper reports a further study into the nature and magnitude of the chemical interaction phenomena and their effect on the age-hardening response of this composite system.

EXPERIMENTAL PROCEDURES

The composites employed in this investigation were based on a heat­treatable matrix of 6061 (Al-l %Mg-0.6%Si-0.5%Cu-0.2%Cr) reinforced with SAFFIL (ICI tradmark) which is a 3f,lm diameter short alumina fibre. These composites were manufactured by a pressure infiltration method/I 1/ using 0.26 volume-fraction fibre preforms. The experimental details of this process have been described in two earlier papers/l,5/.

All the materials were heat-treated in two stages, (i) solution-treated at 529°C followed by water quenching, and (ii) aged isochronally for 30 minutes at temperatures between 20 and 300°C. The hardness of the materials were measured in the solution-treated condition and also following isochronal ageing. Microscopy was carried out using optical and scanning electron (SEM) microscopes, and the distributions of the alloying elements in the matrices were characterised by microprobe analysis.

RESULTS

The unreinforced matrix alloy exhibited normal age-hardening throughout the heat-treatment st,ages and its isochronal ageing response is shown in fig 1. Peak hardness was developed following 30 minutes ageing at ISOoC, which was identical to the behaviour observed in wrought 6061. In contrast the age-hardening of the composites was extremely variable. The ageing response at any point in the composite showed a similar peak hardening temperature to the unreinforced alloy (fig 1), but the magnitude of the peak hardness varied systematically through the section of the composite. Fig 2 shows this effect for composites peak-aged at lSOoC following solution­treatment for 1 hour. In this material there was little variation in the as-cast hardness through the section of the composite, however when solution treated, and when subsequently peak-aged, a hardness gradient developed. This varied from a normal hardening response at the top of the composite

229

(relative to the infiltration direction) to impaired hardening at the base. The magnitude of this phenomenon was strongly dependent on the solutionising time, as can be seen from fig 3. As the solutionising time was increased from 1 to 8 hours, impaired hardening developed throughout the composite.

The absence of a significant hardness gradient in the as-cast condition suggests that this phenomenon was not due to local variations in the reinforcement volume fraction, and this was confirmed by optical microscopy. This showed no systematic variation in fibre fraction through the thickness of the composite. Microscopy did however, reveal changes in the microstructure close to the base of the composite (within 50Ilm). In this narrow region large second-phase particles of a grey constituent were present in the as-cast, and heat-treated conditions. Areas containing these particles showed hardness values above those achieved elsewhere in the composites, and also showed substantial increases in hardness during heat-treatment (fable 1).

The development of hardness gradients following solution-treatment suggests that the observed effects resulted from variations in the solid solution alloy contents. This was confirmed by extensive microprobe analysis. In composites solutionised for one hour these measurements revealed significant segregation of alloying elements to the fibre/matrix interfaces. Digimapping showed a high concentration of Si on all the fibre surfaces and in regions exhibiting impaired hardening, magnesium enrich­ment of the fibre/matrix interfaces. This magnesium segregation was present in the as-cast condition, and was increased by solutionising and subsequent peak-ageing. On extending the solutionising time from 1 to 8 hours more extensive magnesium segregation was observed and this spread to fibre/ matrix interfaces throughout the composite thickness.

Microprobe analysis also identified the microstructural constituents in the base regions of the composites. These were shown to be Si-rich phases, and from their morphologies and analyses were identified as both primary and eutectic silicon particles. Each of these particles was also associated with magnesium segregation at the particle/matrix interface.

DISCUSSION

It is now well established that the age-hardening of 6061 is associated with the formation of the intermediate precipitate pl-Mg2Si/12/. This metastable phase exists over a narrow stoichiometry range, and its formation is therefore impaired by inbalances in either the magnesium or silicon contents of the matrix alloy. Chemical inbalance introduced by interactions can therefore give rise to variations in age-hardening response. This appears to be the source of impaired hardening in the matrices of these MMC. It is well known/10,13/ that o-alumina fibres have silica-rich surfaces which exhibit a strong affinity for magnesium/14/. The phenomena shown in figs 2 and 3 appear to result from such an interaction process.

During composite fabrication the liquid alloy at the inftltration front spends a considerable time in contact with the fibre array. Conversely liquid metal at the top of the preform has a shorter residence time. As a result of interaction between the fibre surfaces and the matrix alloy, significant

230

removal of magnesium can therefore occur. However, because of variations in the liquid/fibre residence times this will result in different levels of segregation between the base of the composite (the original inmtration front) and the top. Since the normal magnesium content of the alloy is low the effect of such a gradient on the as-cast properties is small. However it has a significant effect in both the solutionised and age-hardened conditions. The decreasing amount offree magnesium towards the composite base means that on solutionising smaller amounts enter solid solution. Following short solutionising times this results in a gradient in hardness between the base and the top of the composite, and on subsequent ageing a gradient due to variation in the amount of solid solution magnesium available for ~1-Mg2Si formation. Further evidence for such an interaction process comes from the effect of extended solutionising (fig 3). Figs 2 and 3 show that as the time available for interaction at elevated temperature increases, impaired hardening develops throughout the composite. Microanalysis confirmed that this was associated with increased segregation of magnesium to the fibre/ matrix interfaces. The impaired hardening observed in these composites can therefore be simply explained in terms of interactions between the silica-rich surfaces of the alumina fibres and magnesium in the matrix alloy. Such a mechanism cannot, however explain the increased hardness and superior age­hardening observed in the narrow region at the base of the composites (Table 1).

The region at the base of the composites was characterised microstructurally by the presence of primary silicon which produced elevated hardness in the as-cast condition. The presence of this phase showed that these regions had silicon contents well in excess of the solid solubility of Si in AI. This is strange since the silicon content of 6061 is low. However, in addition to the alloy component, silicon was also present as Si02 in both the fibre and the binder (which imparts handling strength to the fibre preform). The presence of this phase therefore suggests that extra silicon must have entered solid solution via the reaction:

3 Si02 + 4 Al -+ 2Alz03 + 3Si which is thermodynamically possible under any of the conditions present during fabrication. The location of primary Si at the base of the composites suggests that Si02 reduction must have occurred most strongly at the inmtration front, probably as the liquid alloy made contact with the preform binder. Microanalysis showed that these Si particles were also associated with magnesium segregation. However, the product formed was quite distinct from that observed on the fibre surfaces, and is believed to be Mg2Si. Re-solution of this reaction product was possible and heat-treatment gave the age-hardening effects in table 1. The hardness values developed in such matrix regions were in excess of those observed in the normal unreinforced alloy. This is most likely due to the effect of the high silicon content which is known to enhance the density of GP zones, and therefore the hardening provided by the ~lMg2Si /15/.

CONCLUSIONS

8-alumina fibre arrays degrade the age-hardening characteristics of

231

aluminium alloy 6061 matrices. This degradation arises from subtle chemical interactions between the fibre surfaces and active matrix alloying elements. These interactions remove magnesium from the matrix alloy and therefore suppress the formation of the age-hardening ~1 intermediate precipitates.

In preform infiltrated composites the matrix alloy also reacts with other microstructural constituents, including the preform binder. In silica bound preforms this produces high silicon concentrations and can initiate further interactions with matrix alloying elements. Such interactions also lead to anomalous age-hardening.

REFERENCES

1 - Friend C.M., Journal of Materials Science, 22 (1987) 3005-3010. 2 - Friend C.M., Proc. ICCM6 (p. Matthews et al eds), (1987) Vol 2,

402-411, Elsevier Applied Science. 3 - Rack H. 1. ,Proc. ICCM6 (F. Matthews et al eds), (1987) Vol 2. 382-

389, Elsevier Applied Science. 4 - Abis S. and Donzelli G., Journal of Materials Science Letters,

7(1989) 51-52. 5 - Friend C. M. and Luxton S. D. - Journal of Materials Science,

23(1988) 3173-3180. 6 - Ceresara S. and Fiorini P., Powder Metallurgy, 1 (1979) 1-4 7 - Ceresara S. and Fiorini P., Powder Metallurgy, 4 (1979) 210-213. 8 - Clyne T. W., Proc ICCM6 (F. Matthews et al eds), (1987) Vol 2, 275-

286, Elsevier Applied Science. 9 - Trumper R., presented at 'Metal Matrix Composites: structure and

property assessment', The Royal Aeronautical Society, London, November 1987.

10 - Dinwoodie 1. and Horsfall!., Proc ICCM6 (F. Matthews et al eds), (1987) Vol 2, 390-401, Elsevier Applied Science.

11 - Clyne T. W., Bader M. G. Cappelman G. R. and Hubert P. J., Journal of Materials Science, 20 (1985) 85-96.

12 - Polmear I. J., 'Metallurgy of Light Metals' (1981) 15, Edward Arnold, London.

13 - Fox S. and Flower H. M. presented at 'Metal Matrix Composites: structure and property assessment', The Royal Aeronautical Society, London, November 1987.

14 - Capelman G. R., Watts J. F., and Clyne T. W., Journal of Materials Science, 20 (1985) 2159-2168.

15 - Ceresara S., Di Russo E., Fiorini P. and Giarda A., Materials Science and Engineering, 5 (1969/70) 220-227.

Table 1. Hardness At Extreme Base of Composite

1 Hour Solutionised 8 Hour Solutionised

As Cast

100 100

Solutionised

188 177

Peak-Aged

193 210

232

120

- 100 >

J: .. :: 80 c ." c; J: 60

40

0 Ageing re""t,atur. ('CI

Hardn.t:,u

Fig 2 Through thickness hardness traverses in a 0.26 Vc composite (solutionised 1 hour)

110

120

1<0

120

TOO

10

60

Fig 1 Isochronal age-hardening characteristics of the unreinforced alloy and 0.26 Vr composite 12.5mm from top of preform

I (n. , Uration + dlfeCl lon

~Hd'··

10

~ 1", . II,ol.gn

"'«Tf ,T'.

T2 "

100 .god ~ •• u. IIO'C Fig 3 Through thickness -===~~~~~~~;;;I~.~'U~'~s~"u~,,~,"~,,~, hardness traverses in a

As · ,"" 0.26 Vc composite (solutionised 8 hours) 10

60

10 " O'SIOI'I(C hOM lop of pt e'O IM, d IMM)

HOT WORKING BEHAVIOR OF DISCONTINUOUS SiC/AL COMPOSITES OBTAINED BY RHEOCASTING

B. COUTAND, F. GIROT, Y. LE PETIT-CORPS·, J.M. QUENISSET

ABSTRACT

Laboratoire de Genie Mecanique - /UT-A de Bordeaux • Laboratoire de Chimie du So/ide

351 cours de /a Liberation - 33405 TALENCE - France

Composites based on the AI-7Si-O.6Mg alloy and reinforced by short SiC (NICALON) fibers were processed by rheocasting giving the matrix a globular microstructure. Tensile tests were performed at high temperatures with various strain rate. The presence of fibers impeds the plastic deformation giving rise to strong strain hardening effect and low strain rate sensitivity compared to the unreinforced matrix deformed in the highest range of temperature. The observation of a stronger activation energy for the composites is correlated with the modification of residual stresses within the matrix.

INTRODUCTION

The discontinuous reinforcement of aluminum alloys (A.M.C.) by moderate volume fractions of short ceramic fibers can be sufficient to give these composites attractive performance for applications requiring for example good stiffness, fatigue, and wear resistance, as well as relatively low thermal expansion coefficient. Unfortunately, the use of conventional metal-working techniques for mass-production is impeded by the higher resistance to deformation and lower ductility of A.M.C. compared to aluminum alloys 11,21. Therefore, defining the hot working conditions allowing high deformation rate with low flow stress and without the occurence of extensive fiber degradation and cavity formation, is a main concern.

234

During the last five years, several studies have contributed to better understand the mechanisms controlling the flow of an aluminum matrix in presence of a stiff and brittle reinforcement. Most of these studies deal with A.M.C. reinforced by particules or whiskers and processed by powder metallurgy /3, 8/. Only few of them concern A.M.C. reinforced by short fibers and elaborated by squeeze casting or rheocasting /9,10/.

The principal goal of these studies was to adjust the hot working procedure of available aluminum alloys to the presence of fibers or particules. In contrast, the objective of the present contribution is to pOint out the effect of a globular microstructure obtained by rheocasting on the mechanical interaction between fiber and matrix, depending on the strain rate and temperature. This aim results in the use for this first approach of an A.M.C. which is not suitable for easily reaching conditions of superplasticity, but which facilitates the examination of the fiber-globule interaction.

1 - EXPERIMENTAL

1.1. Material

The investigated A.M.C. was produced by rheocasting according to the following procedure :

(1) an AI-7Si-0.6 Mg alloy is heated up to complete melting, (700°C),

(2) after degasing, a vigorous stirring of the melt in the liquidus-solidus range of temperature (600°C) leads to a semi solid mixture made of aluminum globules within a silicon enriched liquid,

(3) then, 10 vol. pct. SiC (NICALON) short fibers (about 12 11m in diameter and 6 mm in length) are progressively incorporated within the slurry while stirring the mixture during 20 mn,

(4) after casting, the AMC is cooled under pressure (40 MPa), (5) finally, the A.M.C. is solution treated at 540°C during 4 h,

quenched in water and overaged at 150°C during 8 h. The microstructures of the resulting composites are illustrated

in fig. 1. The fibers are located within the interglobular space and are homogeneously distributed over the whole material. After tension test at room temperature, fractographic analysis reveal very strong fiber-matrix bonding mainly due to the presence of magnesium /11/. As expected, the stiff brittle reinforcement

235

increases the material rigidity of about 12% but reduces very

significantly the ductility (£f", 0.5%).

1.2. Testing procedure

High temperature tensile tests were performed on the specimen illustrated in fig. 2, within the following ranges of :

(1) temperature : 350 - 520°C (2) strain rate : 10-5 - 2 x 10-3 s - 1 Before loading, the specimens were heated at 5°C. mn- 1 and

maintained at the testing temperature during 15 mn.

2 - RESULTS

The general features of the stress strain curves are illustrated in fig. 2. The fibrous reinforcement increases significantly the strain hardening effect giving rise to higher flow stress. On the other hand, a slight strain softening effect can be observed in the case of the AMC for which the specimen strictions are very small. Furthermore,

the fracture strains £f related to the A.M.C. remain three time

smaller than those corresponding to the unreinforced matrix (120% at 520°C).

2.1. Flow stress dependence of strain rate

The flow stress 00 were derived from the 0-£ curves for two different levels of strain related to the domain of strain hardening and steady state flow stress.

Fig. 3 shows the influence of E on 00 depending on the testing

temperature T. The linear feature of the Ig 0o-Ig £ curves allows the

° 0 -£ relationship to be represented by the following equation :

00 = G(£, T) Em (1) where m is the apparent strain rate sensitivity exponent which simultaneously represents the influence of any change in microstructure and the actual strain rate sensit,ivity of the material; m varies from 0.05 to 0.5 for the unreinforced matrix and increases slightly from 0.1 to 0.2 for the A.M.C. depending on the testing temperature and range of strain rate.

2.2. Temperature dependence

236

Plotting Ig 0'0 v.s 1/T as shown in fig. 4 gives rise to thermal activation curves which evidence a strong influence of the temperature. The A.M.C. exhibit a quasi linear dependence allowing the determination of only one value of the activation energy a for each strain rate over the whole range of temperature. However, the curves related to the unreinforced matrix deviates from a linear representation particularly for the smallest strain rates. Nevertheless, attempting. to represent the 0'0 - e relationship by the following relation gives indicative values of a :

0'0 = H(E,e) e ma/RT (2)

For the A.M.C., the value of a is about 260 +/- 15 kJ/mole, while it is ranging from 150 to 190 kJ/mole for the unreinforced matrix.

3 - DISCUSSION AND CONCLUSION

The previous results and micrographic observations lead to the following remarks:

(1) The stronger increase in the flow stress during the strain hardening of the A.M.C. compared to the unreinforced matrix is consistent with the presence of residual stresses related to the mismatch in the thermal expansion coefficient of the matrix and fiber,

(2) in a wide range of temperature both unreinforced matrix and A.M.C. exhibit low strain rate sensitivities as expected for this type of alloy,

(3) the very significant increase in the strain rate sensitivity of the unreinforced matrix deformed at the highest temperatures and lower range of strain rate is related to the evolution of the microstructure by precipitate dissolution within the globules,

(4) the lack of such an effect in the case of the A.M.C. leads to consider 'that the influence of fibers is predominant on that of precipitates. Also, the ageing treatment could be less effective in the AMC, due to a smaller volume fraction of larger precipitates,

(5) the difference in the activation energies obtained for the unreinforced matrix and A.M.C. shows that the mechanisms of relaxation related to the bulk diffusion in aluminum (",150 kJ/mole ) are significantly accelerated by the high density of residual stresses in the matrix. For instance and as shown by the higher density of cavities in the AMC, the compression of the fiber-matrix interface is released as the temperature is increased which

237

facilitates the built up and coalescence of vacancies, (6) although the flow stress and as a consequence the fiber

shortening remain very significant at 520°C, the fibers are still long enough compared to the length required for load transfer,

(7) as a result the hot working of the A.M.C. is expectable provided to enable the fibers to move within the interglobular space that is allowing globule sliding rather than deforming. Since the eutectic volume fraction is small, the flow stress within the eutectic must be so low, that the alloy temperature has to be adjusted close to the eutectic temperature.

REFERENCES

1- F. Girot, J.M. Quenisset and R. Naslain, Composite Sciences and Technology, Vol. 30, N°3, (1987)155-184

2- F. Girot, L. Albingre, J.M. Quenisset and R. Naslain, J. of Metal, Nov. (1987)18-21

3- M.Y. Wu and 0.0. Sherby, Scripta. Met. , 18 (1984) 773-776 4- B. Derby, Scripta. Met., 19 (1985) 703-707 5- P. Jarry, W. Loue and J. Bouvaist, Proceeding of 'ICCM VI, F.L.

Matthew, N.C.R. Buskell, J.M. Hodkinson and J. morton (eds) Elsevier, London (1987) 2350-2361

6- M. W. Mahoney and A.K. Ghosh, Metall. Trans. A, 18A (1987) 653-661

7- J.R. Pickens, T.J. Langan, R.O. England and M. Liebson, Metall. Trans.A, 18A( 1987) 303-312

8- T.G. Nieh, Metall. Trans 15A (1984) 139-146 9- C.A. Stanford-Beale and T.w. Clyne, Proceedings of the 9 t h

RISO International Symposium on Metallurgy and Materials Science, S.1. Anderson, H. Lilholt, O.B. Pederson,(eds) Roskilde, denmark, 5-9 sept.(1988) 479-484

10- M. Suery, C. Milliere, in "Advanced Materials Research and Developments for Transport", P. Lamicq et al. (eds.), (1985) 241-248

11- F. Girot, J. M. Quenisset, R. Naslain, B. Coutand and T. Macke Proc. ICCM VI ECCM II, London, F.L. Mattew et als. (eds), 2 (1987) 330-339

ACKNOWLEDGMENT The authors acknowledge the contribution of Dr. D. Spataro, and

the assistance of L. Albingre and G. Escusa in manufacturing and machining the composites.

238

ca a. :E f/) f/) (I) 10. -en (I) ::::J 10.

I-

a 100 ).lm

L--...J

b

Fig 1 : Optical micrographs of the microstructure a. Unreinforced AI 7Si O.6Mg matrix b. AI 7Si O.SMg reinforced by 10% SiC (N)

20

b

a 10 If 0

C")

T= 520°C . E= 2.10-3 s-1

0 0 2 4 & 8 10

True Strain, %

Fig 2 : True stress - true strain curves for a. unreinforced AI 7Si O.SMg matrix b. AI 7Si O.SMg reinforced by 10% SiC (N)

12

7

~ :E ~

en ffi a: t;

~ ...I I&.

100

10

1 10.5

239

100

III --Il. :E

~ w 10 a:

• 350°C t; iii 350°C • 400 °C 3: • 400°C • 450 °C 0 • 450°C ...I

• 500 °C I&. • 500°C

· 520 °C · 520°C

10-4 10.3 10.2 1 10-4 10.3 10.5

STRAIN RATE I 8-1 STRAIN RATE I 8-1

Fig 3 : Influence of the strain rate on the flow stress a. Unreinforced AI 7Si 0.6Mg matrix b. AI 7Si 0.6Mg reinforced by 10% SiC (N)

10.2

100~----------------~ 100~----------------~

10 ~ /I:~ • '0-5.·,

• 5.10-5 s-1 a 3.10-4 s-1 • 2.10-3 s-1

10

• 10-5 s-1 • 5.10-5 s-1 a 3.10-4 s-1 • 2.10-3 s-1

'+-~~~~~~T=~ 1,2 1,3 1 ,4 1,5 1 ,6 1,7 1,2 1,3 1,4 1,5 1,6 1,7

11T I 1~3 °K-1 11T I 10-3 °K-1

Fig 4 : Influence of the temperature on the flow stress a. Unreinforced AI 7Si 0.6Mg matrix b. AI 7Si 0.6Mg reinforced by 10% SiC (N)

INTERFACES

Chairmen: Dr A. K. DHINGRA E.!. Du Pont de Nemours Pr R. NASLAIN Laboratoire des Composites Thermostructuraux

ETUDE DE L'INTERFACE FIBRE·ELASTOMERE DANS LES COMPOSITES A MONOFILAMENT

M. NARDIN, E.M. ASLOUN, M. BROGLY, J. SCHULTZ

Centre de Recherches sur /a Physico-Chimie des Surfaces SOlides C.N.R.S. - 24 Av. du President Kennedy - 68200 MULHOUSE - France

ABSTRACT

For the oxidized carbon fibre-EVA system, the existence of an interphase layer between the fibre and the matrix, having an elastic modulus close to the one of the elastomer in its glassy state, can explain the results obtained by fragmentation on single fibre composites. In the present study, viscoelastic measurements on unidirectional composites confirm the existence of such a layer. Hence, the mechanical (average modulus of elasticity) and physical (thickness) properties of this interphase are determined by means of different theoretical approaches. It appears that all the theoretical analyses are in disagreement with experimental observations.

1- INTRODUcnON

La resistance au cisaillement a l'interface fibre-matrice est l'un des parametres les plus importants influen~ant les proprietes mecaniques des materiaux composites. Sa valeur depend de toute modification affectant les proprietes de l'interface fibre-mattice. Le test de fragmentation de la fibre dans un composite a renfort monofilamentaire est tres bien adapte pour mesurer directement cette resistance au cisaillement interfacial. Une charge en traction, appliquee a l'eprouveue dans l'axe de la fibre, est transmise de la mattice a la fibre qui se fragmente jusqu'a ce qu'une longueur critique lc des fragments soit atteinte. La determination des longueurs des fragments permet ainsi d'obtenir une valeur de Ie et de calculer en suite la resistance au cisaillement interfacial. La valeur meme de cette longueur critique peut etre consideree comme une bonne mesure du transfert de charge de la matrice veTS la fibre. En effet, selon les equations analytiques provenant du modele de Cox /1/, Ie facteur de forme critique lr/d, d etant Ie diametre de la fibre, devrait varier, en premiere approximation, comme (Ef/Em)ll2, racine carree du rapport des modules d'elasticite de la fibre et de la matrice, cela dans Ie cas d'une tres bonne adhesion fibre-mattice.

Une etude recente /2/ a montre que l'analyse de Cox etait verifiee dans un large domaine de valeurs du rapport EflEm, queUe que soit la nature des fibres et des matrices

244

utilisees. Ainsi, en coordonnees logarithrniques, la variation de ldd en fonction de Ef/Em est une droite de pente 1/2, tres proche de celie theoriquernent prevue. Cependant, il a ete observe que les resultats correspondant aux matrices 6lastomeres, notamment pour des temperatures (T) superieures a leur temperature de transition vitreuse (T g)' s'ecartaient notablement de la droite maitre sse proven ant du modele de Cox, les valeurs experirnentales de ldd etant alors tres inf6rieures a celles que prevoit l'analyse theorique. Ce phenomene est particulierement clair dans Ie cas d'un systeme fibre de carbone oxydee/matrice poly(ethylene-acetate de vinyle) (EVA), comme Ie montre la figure 1, pour des temperatures superieures a la temperature de transition vitreuse de l'EV A, egale a environ -36°C.

Pour une valeur donnee du rapport !=If/Em, tout se passe donc comme si Ie transfert de charge de la matrice a la fibre etait plus efficace dans le cas d'une matrice elastomere que pour une autre matrice (therrnoplastique ou therrnodurcissable), puisque la longueur critique lc est alors plus faible. Trois phenomenes peuvent etre invoques pour expliquer un tel comportement: (i) un effet de friction dii a la contraction transversale de l'elastomere lors du test de fragmentation, (ii) un effet viscoelastique deja rencontre par ailleurs /3,4/ dans les phenomenes d'adhesion, (iii) l'existence d'une couche a !'interface fibre-matrice, ou "interphase", dont les propri6tes mecaniques different notablement des proprietes massiques de la matrice. n a ete montre /2/, par des mesures de lc a differents aliongements relatifs, que l'effet de friction ne pouvait pas etre tenu pour responsable du phenomene observe. Si un effet viscoelastique est mis en jeu, un facteur multiplicatif dependant de la vitesse de sollicitation et de la temperature devrait alors etre pris en compte. L'existence d'une couche interfaciale, cependant, semble etre l'hypothese la plus prometteuse /2/. Ainsi, au cours de cette etude, nous nous proposons, dans un premier temps, de mettre en evidence l'existence d'une tene interphase a l'aide de mesures de viscoelasticite et, dans un deuxieme temps, d'estimer theoriquement les proprietes physiques et mecaniques de cette demiere.

2- CONDIDONS EXPERIMENT ALES

Des fibres de carbone T300 (Torayca) de haute resistance, non traitees d'une part et ayant subi une oxydation de surface par un traitement electrolytique d'autre part, ont ete utilisees au cours de cette etude. Un seul type de matrice elastomere, un poly(ethylene-acetate de vinyle) (EVA) (DuPont, Elvax 150) a toujours ete employe.

Les conditions experimentales concernant Ie test de fragmentation a differentes temperatures, sur les composites a monofilament fibre de carbone oxydee-EV A, sont decrites dans une etude precedente /2/.

Les mesures de viscoelasticite ont ete effectuees a l'aide d'un appareil Metravib MAK 03 sur des composites EVA renforces unidirectionnellement a 0, 5, 15, 30% en poids par des fibres de carbone vierges ou oxydees. Les echantillons ont ete sollicites en traction-compression perpendiculairement a l'orientation des fibres, a des temperatures comprises entre -100 et + 100°C et a des frequences variant de 7,8 a 125Hz.

De surcrolt, une analyse par enthalpimetrie differentielle a balayage de temperature (DSC) a ete effectuee sur ces differents composites au moyen d'un calorimetre Mettler TA 3000 dans une gamme de temperature comprise entre -150 et +150°C et a une vitesse de montee en temperature constante, egale a lOoC/min.

3- RESULTATS

L'influence du taux de fibres a d'abord ete analysee par une etude isochrone (7,8Hz) en viscoelasticimetrie. Nous nous sommes ainsi interesses aux variations du module de perte E" des materiaux en fonction de la temperature. Ces variations presentent des pics de relaxation au voisinage de la transition vitreuse de la matrice EVA,

245

soit environ -40°C. Les valeurs de la temperature T m' correspondant au maximum des pics dans les differents cas, sont regroupees dans Ie tableau I. On cons tate que pour les fibres de carbone non traitees, T m ne varie pas en fonction de la teneur en fibres, alors que, pour les fibres oxydees, T croit lineairement avec cette derniere. Une telle variation indique donc que, dans Te cas des fibres oxydees, la relaxation concernee est affectee par la presence et la quantite de fibres au sein du materiau. L'augmentation de T m traduit une moindre mobilite des chaines polymeres au voisinage des fibres. 11 apparait ainsi que la nature de surface des fibres, dont dependent les interactions physico-chimiques fibre-matrice, est un facteur essentiel.

Pour une teneur en fibres constante, egale a 15% en poids, les variations de la tangente de l'angle de perte (tgB) en fonction de la temperature et de la frt5quence ont ete ensuite etudiees. 11 est possible ainsi de determiner les temps de relaxation 't correspondant aux temperatures maximales Tm des pics observes. La loi d'Arrhenius reliant 1: a l'energie apparente d'activation Ea de la relaxation:

't = 'to exp( EafRT m) ou Rest la constante des gaz parfaits, peut etre alors consideree comme applicable dans Ie domaine restreint de temperatures ou cette relaxation a lieu.

La figure 2 indique clairement que la loi d'Arrhenius est applicable pour la matrice EVA et pour les composites con tenant des fibres vierges ou oxydees. Les valeurs calculees de Ea sont donnees dans Ie tableau II. On constate que l'energie apparente d'activation est identique pour la matrice seule et pour Ie systeme contenant des fibres non traitees. Cette energie, par contre, est environ deux fois plus grande dans Ie cas du composite contenant des fibres oxydees. Ces resultats confrrment ainsi les precedentes conclusions, a savoir l'existence dans les composites a matrice EVA renforcee par des fibres de carbone oxydees d'une couche de polymere de moindre mobilite au voisinage des fibres. Notons, cependant, qu'aucune difference ne peut etre decelee entre les thermogrammes DSC des differents produits, queUes que soient la teneur et la nature de surface des fibres.

4- DISCUSSION

Apres avoir mis en evidence l'existence d'une couche interfaciale au voisinage des fibres de carbone oxydees dans une matrice EVA, il est interessant main tenant d'essayer d'estimer les proprietes physiques et mecaniques d'une tene interphase et notamment son epaisseur et son module d'elasticite moyen. Notons tout d'abord que dans une etude precedente /2/, il a ete suppose que Ie comportement mecanique d'une teUe couche etait equivalent a celui de l'elastomere dans son etat vitreux, queUe que soit la temperature. Si tel est Ie cas, il est possible de determiner Ie module d'elasticite moyen Em * de cette interphase en extrapolant a T>T.,g les valeurs du module d'elasticite Em de la matrice mesure pour des temperatures interieures aT g' Le tableau III montre que Em * est de 20 a 60 fois superieur environ a Em. n apparait aussi immediatement (figure 1) que, pour des temperatures superieures a T g et en employant Em * au lieu de Em' les valeurs experimentales du facteur de forme critique ldd se replacent sur une droite de pente 1/2, en accord avec l'analyse 'de Cox /1/, ce qui tend a confirmer la realite d'une telle interphase.

En partant de cette meme hypothese concernant la valeur du module d'elasticite de l'interphase, il est possible de determiner l'ordre de grandeur de l'epaisseur ei de cette derniere. Deux analyses differentes permettent d'effectuer ce calcul. Premierement, l'analyse presentee par Piggott /5,6/, qui est equivalente a ceUe de Cox mais prend en compte l'existence d'une region interfaciale bien definie et dont les proprietes sont independantes de ceUes de la fibre et de la matrice,peut etre developpee dans Ie but d'acceder a l'epaisseur de cette couche. Un raisonnement identique a celui que l'on applique dans Ie modele de Cox pour etablir la relation entre ErlEm et lcfd est retenu /2/.

246

Dans l'analyse de Theocaris n -WI, au contraire, il est suppose que Ie module d'elasticite de l'interphase varie continfiment de Ef it la surface de la fibre it Em it une distance ei de la fibre. La valeur moyenne du module d'elasticite de l'interphase est consideree egale it Em *, precedemment determinee. L'epaisseur ei de la couche interfaciale verifie alors les equations suivantes:

Em* = Ern/2 + Ef{(BTl -B1!2)/(l-2TJ)(l-B)1!2 - BTl!2)

avec 2TJ(1-B1I2) + 1 = Em/EfBTl et B= (rf/(rf+ei»2, TJ etant un coefficient dont l'inverse traduit l'amplitude des interactions fibre-matrice.

Le tableau III presente les valeurs de ei it des temperatures superieures it la temperature de transition vitreuse de l'EV A. Un ecart important peut etre observe entre les resultats obtenus par ces deux approches. L'analyse de Theocaris conduit it des epaisseurs de l'interphase anormalement elevees, tres superieures it l'epaisseur de la matrice elle-meme, egale it 1mm dans tous les cas. Notons, cependant, que la valeur du coefficient TJ reste con stante et egale it 0,73 ± 0,03, queUe que soit la temperature, ce qui laisserait supposer que les interactions fibre-matrice ne sont pas modifiees dans Ie domaine restreint de temperatures etudiees, ce qui est raisonnable.

De surcroit, les plus faibles valeurs de ei, calculees suivant l'analyse de Piggott, sont encore tres largement superieures it celles que 1'0n peut obtenir it partir de l'analyse de Lipatov /11/. Selon cette theorie, l'epaisseur moyenne de la couche interfaciale dans les elastomeres renforces depend du rapport des variations de capacite calorifique ~Cp de la matrice seule et des composites, dans la region de transition vitreuse de l'elastomere. Dans notre cas, les resultats obtenus en DSC conduisent, selon cette analyse, it des valeurs de ei egales it quelques dizaines, voire quelques centaines, de nanometres, que 1'0n soit en presence de fibres non traitees ou de fibres oxydees. Ce dernier point est en desaccord avec les precedents resultats obtenus en viscoelasticimetrie.

Enfin, il faut observer que, selon les approches de Piggott et de Theocaris, les valeurs des epaisseurs de l'interphase augmentent avec la temperature. Ce fait est en total desaccord avec des observations recentes, effectuees sur des elastomeres charges /12/. Ces observations semblent montrer qu'une augmentation de la temperature se traduit plus par une diminution notable de l'epaisseur de la couche interfaciale que par une modification des valeurs de module de cette derniere.

5- CONCLUSIONS

Les resultats de fragmentation sur composites it monofilament fibre de carbone oxydee-EVA laissent supposer l'existence entre la fibre et la matrice d'une couche interfaciale, dont les proprietes sont differentes de celles de la matrice en masse. L'etude presente, it l'aide de mesures de viscoelasticite sur composites unidirectionnels it differents taux de fibres, confume cette hypothese. Une teUe couche existe bien dans la mesure ou les interactions fibre-matrice sont fortes, ce qui est Ie cas avec les fibres de carbone oxydees. Cependant, aucune approche theorique retenue dans cette etude ne peut nous renseigner sans ambiguite sur les proprietes physiques et mecaniques de l'interphase. Il n'est pas simple, en effet, de modeliser cette derniere, puisque de nombreux parametres difficilement mesurables (module moyen, module it la surface de la fibre, epaisseur ... ) doivent etre pris en compte. Le peu de connaissance que nous avons, it l'heure actuelle, sur Ie comportement des polymeres au voisinage des fibres dans les materiaux composites, nous conduit it emettre des hypotheses discutables sur Ies valeurs de teis parametres. Neanmoins, dans l'avenir, des mesures fines de module, par exemple, ou d'autres approches theoriques devraient nous renseigner plus precisement sur l'etat d'une telle interphase.

247

REFERENCES

1- H.L. Cox, Brit. J. Appl. Phys. 3 (1952) 72. 2- El. M. Asloun, M. Nardin et J. Schultz, "Stress transfer in single fibre

composites: effect of adhesion, elastic modulus of fibre and matrix, and polymer chain mobility." a paraitre dans J. Mater. Sci ..

3- A.N. Gent et J. Schultz, J. Adhesion 3 (1972) 281. 4- E.H. Andrews et A.I. Kinloch, Proc. Roy. Soc. A332 (1973) 385. 5- M.R. Piggott, J. Mater. Sci. 13 (1978) 1709. 6- M.R. Piggott, Polym. Compos. 8 (1987) 291. 7- G.C. Papanicolaou, P.S. Theocaris et G.D. Spathis, Colloid Polym. Sci. 258

(1980) 1231.8. P.S. Theocaris, Colloid Polym. Sci. 263 (1985) 863. 9- P.S. Theocaris et T.P. Philippidis, J. Mater. Sci. 22 (1987) 3407. 10- P.S. Theocaris et A.G. Varias, Rheol. Acta 26 (1987) 322. 11-Yu.S. Lipatov, "Physical Chemistry of Filled Polymers", International

Polymer Science and Technology, Monograph n02, transl. from the Russian by R.I. Moseley, Ed. Rubber & Plastic Res. Ass. of Great Britain (1979), p.66.

12-B. Haidar, communication privee (1988).

taux de T (OC) m

fibres (%) non traitees oxydees

0 -44 -44 5 -42 -42

15 -46 -38 30 -42 -

Tableau I: Variation de T , correspondant au maximum des pies de relaxation E"=f(T) a T",T ,men fonetion du taux de fibres.

g

Ea (kJ/rrol)

EVA 227 ± 14 EVA-fibres non traitees 235 ~ 10 EVA-fibres oxydees 550 - 57

Tableau II: Valeurs de l'energie apparente d'aetivation de la relaxation a TOIT • g

T (OC) E (MPa) E: (MPa) e i (Piggott) e i (Theocaris) m (~) (~)

-20 30.50 524 0 3710 0 12.85 307 27 6070

+20 6.60 179 124 11110 +40 1.67 104 336 15140

Tableau III: Caraeteristiques de l'interphase entre une fibre de carbone oxydee et une matriee EVA en fonetion de la terrperature.

248

1 10 3 cl d

" ..... '~·:40 .--...

-:~----~---:no

Efl Efl E ou E-m m

10 10 10

Figure 1: lc/d en fonction de Ef/E ou Ef/E~ pour un systeme fibre de carbone oxydee-E~ (cerc1es: ca1cu1s effectues avec Ef/E ; carres: ca1cu1s effectues avec Ef/E-; 1es nombres ~res des points experimentaux corres­pond~nt aux temperatures en °e).

o r------r------r------r----~

-2

~4

3,9 '4,0 4,1

Figure 2: Variation de In-[ en fonction de liT : (1) EVA; (2)EVA-fibres non traitees (15%); (3)EVA-fibr~s oxydees (15%) •

GREFFAGE ELECTROCHIMIQUE DE FONCTIONS AMINEES EN SURFACE DE FIBRES DE CARBONE: EFFET SUR LA TENACITE D'UN

COMPOSITE CARBONE·EPOXY

ABSTRACT

B. BARBIER, M, VILATTE., G. DESARMOT

ONERAIOM BP 72 - 92322 CHA TIL LON CEDEX - France

*AEROSPATIALE 15 Rue Pasteur - 92150 SURESNES - France

Electrochemical grafting of amino groups at the surface of carbon fibres (high strength and high modulus) is highly described. Amino groups are covalent bonded to carbon atoms forming a thin and dense surface layer. They are able to react, for instance. with epoxy groups of an organic matrix. We compare toughnesses of composite materials containing untreated or treated (oxidised, aminated) carbon fibres. It is shown that aminating surface treatments are a new way to improve brittle matrix composite material toughness or to get a better adhesive matching between matrix and fibres.

TNTRODUCTION

Le greffage electrochimique de groupements amines a la surface de Ibres de carbone (M. Sanchez /1/, Demandes de Brevet Fran~ais /3.4/. I. Barbier /5/, B. Barbier & coli. /6/) est un procede de tra:itement de mrface qui. a I'inverse des traitements oxydants classiques, a I'avantage lue I'on peut faire varier tres largement la nature - voire I'epaisseur­Ie la couche superficielle de greffons amines. Cela rend a priori possible me action sur la tenacite de materiaux composites carbone-resine,

notamment ceux comportant une matrice fragile ou une matrice ayant une faible adhesion (type PPS) sur des fibres oxydees en surface.

I - METHODE DE GREFFAGE DE SURFACES DE FIBRES DE CARBONE

1.1. Determination des conditions electrochimiques du greffage

Le greffage est obtenu en oxydant electrochimiquement une amine en solution organique dans une cellule dont l'anode est une meche de fibres de carbone (traitement potentiostatique). Pour une amine primaire, la reaction d'oxydation s'ecrit :

250

-e- +. R-NHz --) R-NHz

Le produit forme est un radical cation reagissant avec Ie carbone des fibres (§ 1.3). Un voltamogramme permet de connaltre Ie potentiel de pic d'oxydation Ep de I'amine. Pour traiter des fibres, il est necessaire de se placer a un potentiel de travail VT superieur au potentiel de pic Ep, tout en demeurant en dessous du potentiel de decomposition de la solution. Nous utiliserons l'ethyhinediamine NHz-(CHz)z-NHz (EDA ; Ep = +1,3 volt/ECS) et la triethyIenetetramine NHz-(CHz)z-NH-(CHz)z-NH­(CHz)z-NHz (TETA ; EpNB = +0,97 volt/ECS - EpNBz = +1,25 volt/ECS). Ces conditions valent pour la cellule de voltaml!trie qui a servi pour la determination des conditions operatoires des traitements de surface (avec un fragment de meche de 4 cm de longueur) et pour la cellule pilote dans laquelle on traite de grandes longueurs de fil pour Ia fallti.cation d'eprouvettes composites (§ 2).

1.2. Caracterisation des surfaces traitees : analyses de surface

On a utilise les SIMS et l'ESCA. A partir de fibres haute resistance COURT AULDS XAU non traitees, on a represente :

- en figure I, une analyse ESCA effectuee sur ces fibres XAU non traitees. On ne detecte pratiquement que du carbone, de l'oxygene et des traces d'azote (pies CiS, 01S et N1S) ;

- en figure 2, une analyse ESCA pour un fragment de meche traite dans la cellule de voltametrie avec de l'EDA pendant 5 minutes (VT = +1,6 volt/ECS ; solvant : CHaCN ; electrolyte support (C4Hg)4NBF4 : 0,2 Mil ; concentration d'EDA : 0,2 M/l). On observe un pic N1S tres intense, attestant la presence d'une forte concentration d'azote Ii la surface des fibres. La forme du pic CiS est modifiee apres traitement ; il comporte alors une composante correspondant Ii des liaisons covalentes C-N etablies entre Ie carbone de la fibre et l'amine.

- en figure 3, une analyse en SIMS negative des fibres XAU non traitees, dans des conditions guasi-statigues ;

- en figure 4, une analyse en SIMS negative des fibres XAU traitees. Le pic Ii la masse reduite M/Z = 26 (ions CN-) est devenu beaucoup plus intense que Ie pic Ii la masse 24 (ions CC-). Si on admet que ce der­nier represente Ie. substrat carbone et qu'on puisse s'en servir comme reference, l'extreme voisinage de la surface (epaisseur ::: 0,5 nm) est considerablement plus riche en azote apres traitement : Ie pic Ii M/Z = 26 (ions CN-) devient en proportion du pic Ii la masse 24 de 0,87 a 9,5 fois plus intense. L'augmentation de la teneur en oxygene est moins marquee.

Par consequent, Ie traitement apporte beaucoup d'azote en surface et un peu d'oxygene, les fibres non traitees comportant deja en surface une proportion notable de cet element. La figure 5 est un profil ionigue SIMS obtenu sur un monofilament de carbone : Ie diametre du faisceau primaire d'ions Cs· est 1,4 ~m. La vitesse d'erosion est environ 2 nm/mn. On observe une tres forte diminution du signal eN-, montrant que l'azo-

251

te est tres superficiel. C'est g(meralement Ie cas pour tous les traite­ments. En revanche, la repartition de l'oxygene depend des conditions operatoires et peut presenter des gradients plus ou moins marques.

Le tableau I regroupe les resultats d 'analyses de surface effectuees sur des echantillons traites dans des conditions differentes :

- sur fragment de meche : effet de la temperature et du potentiel. Une augmentation de l'un ou l'autre de ces parametres accelere la cinetique de greffage ;

- sur: fil trait~ en continu : mesures sur les fibres externes et mesures sur les fibres internes dans une meche. On cons tate que Ie traitement est applicable a une installation en continu et qu'U est homogene.

On a egalement reporte les resultats relatifs a une fibre COURT AULDS haut module HMU, avant et apres traitement. Pour tous les traitements, les concentrations atomiques en azote (ESCA) sont comparables ; c'est Ie cas egalement pour d'autres amines. On a donc, passee une certaine duree de traitement, saturation de la surface en azote, I'epaisseur de la couche riche en azote avoisinant Ie nm. 1.3. Mecanisme reactionnel

A titre d'exemple, on donne ci-apres Ie mecanisme de greffage de I'EDA. II est discute en detail en 15/. II Y a deux mecanismes possibles

Oxyda- Diradical-tion cation

Greffage

-e- +. NH2-C2H4-NH2 --) NH2-C2H4-NH2

Radical-cation

-2Hf --)

Deproto­nation

-H' --)

Ces deux formes de greffage se produisent simultanement. La TETA presente de multiples possibilites de greffage analogues a celles de l'EDA.

II - TENACITE ET TENUE EN CISAILLEMENT DE COMPOSITES COURTAULDS XAU/NARMCO 5208

La presence de groupements amines a la surface des fibres permet l'accrochage chimique d'une resine epoxyde selon Ie schema simplifie :

CH2 -CH"""- --) , 0/

252

La nature, les fonctionnalites et la longueur du groupement R permettent la constitution d'interfaces ou d'interphases ayant - par rapport a un traitement oxydant classique - une tenacite differente.

Nous avons mesure la tenacite de composites en mode I ({mergie de creation de surface en mode I ou G1e) par la methode de la double poutre encastree (voir C. Marais & colI. /71 et D. Guedra & colI. 18/). Les echantillons comportent 24 plis ; Us sont prepares a partir de fils traites dans la cellule pilote. Les materiaux de reference sont :

- un composite COURTAULDS XAU (fibres non traitees)/resine NARMCO 5208 ;

- un composite COURTAULDS XAS (traitement oxydant commercial)1 resine NARMCO 5208.

Les fils de carbone ont circule sans courant dans la cellule pilote contenant une solution, afin que l'on retrouve Ie meme enchevetrement que pour les fibres XAU traitees avec EDA et TETA - du reste tres faible mais ayant une action sensible sur G1e.

La figure 6 montre de maniere simplifiee les variations de G1e en fonction de la longueur de propagation a de la fissure dans l'eprouvette. La forme de ces courbes est discutee en 161 et /71. La croissance initiale est due a la formation de ponts de fibres entre les deux levres de la fissure. Au-dela d'une certaine valeur de a, G1e atteint une valeur stabilisee : la quantite de ponts de fibres qui se forme est egale a la quantite de ponts de fibres qui se rompt. Dans cette experience, la mesure de G1e depend de la geometrie de l'eprouvette ; elle ne donne pas une valeur intrinseque de la t{macite du materiau. Mais comme ~ seul parametre ill!! distingue les materiaux entre eux est Ie traitement de surface, on peut etablir une echelle des tenacites. On constate que :

- Ie traitement commercial (XAS/5208) abaisse la tenacite par rapport a l'absence de traitement (XAU/5208). Le traitement commercial conduit a une propagation instable de la fissure, que nous avons schematisee par une ligne brisee refletant la realite ;

- les traitements a l'EDA et a la TETA ne conduisent pas a une valeur de G1e qui se soit stabilisee (elle serait atteinte pour des longueurs de fissure plus elevees) ;

- Ie traitement a la TETA produit une tenacite environ deux fois plus elevee que les autres traitements, en particulier Ie traitement oxydant, avec une propagation stable de la, fissure.

Des mesures de cisaillement interlaminaire (CCIL) et de contraintes de de cohesion fibre-matrice Td obtenue avec Ie test d'arrachement d'un monofilament (pull-out test, voir G. Desarmot & M. Sanchez 191 et M.R. Pigott & colI. /IO/) sont reportees dans Ie tableau II en regard des mesures de G1e.

D'une maniere generale, l'action d'un traitement oxydant est d'aug­menter I'adhesion fibre-matrice et d'abaisser la t(macite du materiau. Le

253

traitement a la TETA augmente simultanement l'adhesion fibre-matrice et la tenacite du materiau composite (c'est egalement Ie cas pour EDA, mais a degre moindre). L'origine de ce phenomime doit etre recherchee au niveau de la tenacite de l'interface, dont l'influence sur la tenacite globale du materiau doit etre beaucoup plus importante qu'il n'y parais­saito La moindre fragilite d'une "interface aminee" est une propriete qui n'a, pour l'instant, pas re~u d'explication precise. On peut penser que la structure de la couche azotee au voisinage de la surface des fibres et les possibilites de mouvement des chaines aminees Ii l'interface favorisent tine augmentation de la tenacite interfaciale.

CONCLUSION

Dans ce bref expose, nous mettons en evidence l'interet du procilde de traitement de surface de fibres de carbone par des substances aminees en milieu organique. Opere dans de bonnes conditions (potentiel, temperature), Ie traitement des fils est homogene et rapide (quelques dizaines de secondes). II est applicable Ii d'autres substances carbonees que les fibres. Dans Ie domaine des composites Ii matrice organique, la variete des substances greffables autorise une veritable adaptation des surfaces traitees aux matrices d'impregnation, pour ameliorer la tenacite des materiaux sans que cela soit aux depens de la tenue en cisaillement, au moins dans une large mesure.

BIBLIOGRAPHIE

1. M. Sanchez, "Traitements de surface de fibres de carbone : effets sur l'adhesion fibre-matrice", These de Doctorat de l'Universite Paris VI, 22 mai 1986 2. M. Sanchez, G. Desarmot, Comptes Rendus des Cinquiemes Journees Nationales sur les Composites (JNC5), Paris, 9-11 septembre 1986, p. 471-487, Pluralis 3. Demande de Brevet Fran~ais n° 84.07814 4. Demande de Brevet Franr;ais n° 86.16841 5. B. Barbier, These de Doctorat de l'Universite Paris VII, Ii paraitre 6. B. Barbier, M. Villatte, M. Sanchez, G. Desarmot, Comptes Rendus des Sixiemes Journees Nationales sur les Composites (JNC6), Paris, 11-13 octobre 1988, p. 115-130, Pluralis 7. C. Marais, M.C. Merienne, P. Sigety, Comptes Rendus des Cinquiemes Journees Nationales sur les Composites (JNC5). Paris 9-11 septembre 1986, p. 3-16, Pluralis 8. D. Guedra, D. Lang, J. Rouchon, C. Marais, P. Sigety, Sixth International Conference on Composite Materials (ICCM6), London, 20-24 July 1987, p. 3.346-3.357, Elsevier 9. G. Desarmot, M. Sanchez, Comptes Rendus des Quatriemes Journees Nationales sur les Composites (JNC4), Paris, 11-13 septembre 1984, p. 449-460, Pluralis

254

Tableau! - Analyses ESCA

Cellule Fibre Poten tiel Dur~e Temp.- \C \N \0 VT traitement rature

nu - Non trait~e - 88 2,2 9,8

nu +1,6 volt/ECS 300 s 20°C 64,1 22,4 12,9

XAU +1,6 volt/ECS 45 s 50°C 63,S 23,8 12 ,1 VOLTA

XAU +2,0 yolt/ECS 45 s 20'C 61.8 25,2 13,0

EIIU +2,0 voH/ECS 90 s 20°C 62,2 21.3 10 ,5

XAU +1,6 voH/ECS 300 s 20·C PILOTE Fibres externes 65,9 23.8 10,3

Fibres internes 65,3 25.8 8.9

Solvant . CH,CN - Elec trolyte ' ethylenediamine (EDA)

ESCF=! 1---~ 5UAf"""~1 RISER SIMS I--I.t ·"

CC n ( .... o-S . ' 1I: - -1 , 1 11 CN -

1 3 ~ . " .- " ,

c

.. I ~ " , c ° u ~

~, "

u

. I I ... . ..

C amu 1 '" t1.aS"s "1)0 lOO

BINDING E:NERCY (cV)

ESCF=! I-- ern 5IJIVl'ICu 1 R!8~R S I MS 1---.- eN"

2 ..... 4 u . , c

8 c ..... 0" g

Jf~ u

. ,~ . ..

(alPlLl ] '" Mass ______ '~ ___ "!B~I~N~~ll~~G~E~N~E~R~G~Y~·~~.~Vl) ____________ J_----------____ ~~~~~ ________________ ~

t;I'It1I:Cill IH1IJI'" ~TM ~u.~ z ..... n' •• I' 1.'·r------------'·"-·..,-=·..! .. ""-~----

5

cc-

19. 1: Spectre ESCA pour fIbres XAU non tra lt ees 1&. 2 : Spectre ESCA pour f ibres XAU trallee. a I'EDA Ig. 3 : Spectre SIMS- pour fibres XAU non tralt~ e. If : ~ : ~~~flir~!~~,,!s; J:.°::,romfr:m~ttU xt..:tlttr":.t: I'EDA

Tableau!! - Mesures mecaniques

Systeme G1C ceIL Td (J/m2 ) (MPa) (MPa)

XAU non traitee/5208 400 69 ± 5 88,5 ± 5

XAS/5208 315 115 ± 10 ::: 125

XAU (EDA)/5208 385* 87 ± 9 117 ± 5

XAU (TETA)/5208 640* 103 ± 9 107 ± 3,5

• Stabilisation non atteinte

Fig, 6: G1C, imergie de creation de surface en mode I (tenacite) de composites XA /5208

600 ~(TETA)/520a 500

400

XAU(EDA)/5208

300 XAS/5208

200

resine

100 ,

5208 (24 plis)

longueur de la fissure

255

ELABORATION EN CONTINU D'UN DEPOT MINCE DE CARBURE REFRACTAIRE EN SURFACE DES FIBRES DE CARBONE:

ABSTRACT

CARACTERISATION DE LA FIBRE CISiC

H. VINCENT, C. VINCENT, J.L. PONTHENIER, H. MOURICHOUX, J. BOUIX

Universite Claude Bemard Laboratoire de Physico-Chimie Minerale I

43 boulevard du 11 novembre 1918 - 69622 VILLEURBANNE - France

Reactive chemical vapour deposition (RCVD) has been used for the coating of the single filaments of a PAN-based carbon yarn by SiC. A continuous process has been developped, it consists to react a mixture SiCI4-H2 with the fiber, at T 1000oC, under. atmospheric pressure. By the aiH of thermodynamic calculations and experimentation, the conditions of RCVD (gaz composition, temperature, reaction time) are opt imized to e labore a coat i ng without degrading the fiber. It is shown that a SiC-B coating, 0.05 micron thick, modified slighty the characteristics and is sufficient to reduce the rate of combustion of the fibers by a factor closed to 100.

INTRODUCTION

Divers composes refractaires (SiC, TiC, B4C, TiB2' TiN .•• ) ont ete proposes pour revetir la surface des fibres de carbone avant leur incorporation dans une matrice metallique, tel que l'aluminium. La presence de ce revetement a pour but de limiter la reaction chimique entre Ie carbone et l'aluminium qui se produit des 500°C /1-2/, soit a une temperature inferieure a la temperature d'elaboration du materiau. Le procede de traitement des fibres doit permettre Ie revetement uniforme de tous les filaments unitaires constituant la fibre, et Ie controle de son epaisseur qui doit rester suffisamment fine pour ne pas affecter les proprietes mecaniques de la fibre initiale.

Dans cette publication, nous montrons que ces dernieres exigences peuvent etre respectees par Ie procede de RCVD (Reactive Chemical Vapour Deposition). Nous avons choisi SiC comme exemple de revetement. Ce carbure est en effet stable thermodynamiquement avec Ie carbone et avec 1 'aluminium a une temperature inferieure a 650 0 c ou avec un

258

alliage Al-Si a une temperature superleure /3/ et SiC doit ameliorer la mouillabilite des fibres /4,5/. Nous montrons que l'elaboration d'un depot mince de SiC peut etre realisee en continu, a la pression atmospherique. L'optimisation des conditions de la RCVD a ete faite en relation avec des tests de rupture en traction sur monofilaments et des tests d'oxydation a l'air.

I-PROCEDE DE REVETEMENT: RCVD

1.1. Principe

Le procede de RCVD se distingue des procedes de CVD classiques par le fait que la phase gazeuse reagissante ne contient pas de carbone, mais seulement un melange d'halogenure et d'hydrogene. La reaction de depot s'ecrit:

SiC1 4(g) + 2 H2(g) + C(s) ~ SiC(s) + 4 HC1(g) (1)

L'interet de ce procede reside dans la necessite de carbone pour que la reaction puisse se poursuivre, avantages sur les procedes de CVD classiques:

la presence de d'ou plusieurs

-la peripherie de la fibre de carbone n'est plus privilegiee, l'infiltration a coeur est possible puisque les reactions de surface sont lentes. -le depot est forme uniquement en surface de la fibre, ce qui permet un traitement en continu sans risque de colmatage du reacteur et avec un taux eleve de conversion de SiC1 4 en SiC.

1.2. Aspect thermodynamique

Bien que l'aspect cinetique du processus soit tres important, la quantite de carbone mise au contact de la phase gazeuse diminuant avec le temps, l'etude thermodynamique du systeme SiC14-H2-C permet de prevoir les tendances d'evolution de ce systeme en fonction des conditions imposees. De l'etude complete de ce systeme /6/, nous pouvons tirer un certain nombre de resultats d'interet pratique.

Les graphes de la ~ representent la composition a l'equilibre des principales especes formees en fonction de la quantite initiale Q de carbone a 1700 K, a P = 1 atm, pour deux compositions de la phase SiC1 4-H2• Les resultats sont rapportes a 1 mole de SiC14• La reaction (1) de formation de SiC est donc accompagnee de reactions secondaires conduisant a la formation de composes gazeux du carbone (CH4 ) et du silicium (SiC1 3 , SiC1 2 , SiHC13 •• ). Les conditions optimales de la RCVD sont celles qui favoriseront la reaction (1) au detriment de toutes les autres:

-en presence d'un exces de carbone, SiC est la seule phase condensee -la quantite de carbone engagee sous forme de CH4 est limitee en reduisant la quantite d'hydrogene dans le melange gazeux initial. -un codepot de Si et de SiC est possible lorsque la quantite de carbone est faible (fig.1b). 11 existe cependant des compositions de

259

la phase H2-SiC14 qui interdisE'nt le depot de silicium meme en absence de carbone (~) (partie hachuree de la fig.2). -la dilution du melange initial par un gaz inerte permet d'abaisser la temperature de depot de SiC en gardant des taux de conversion de SiCl en SiC et de C en SiC proche de l'unite. La ~ montre l'intluence d'une dilution par l'argon sur ces taux pour un melange initial caracterise par 2,3 moles de H2 pour 1 mole de SiC1 4•

II-RESULTATS EXPERIMENTAUX

2.1. Dispositif

Le reacteur de RCVD a ete decrit precedemment /7!. Ii comprend essentiellement une chambre de depot cylindrique en silice et deux poulies en acier inoxydable qui assurent le defi lement de la fibre et sa mise sous tension electrique. La duree du traitement test assimilee au temps de sejour de la fibre dans la chambre. Un dispositif de distribution de gaz permet toutes les compositions possibles H?_-SiC1 4-Ar. Par la suite, nous avOns dHini le melange initial par ie rapport R entre les debits d'hydrogene et de SiC1 4 et par le debit total DT• La pression totale est la pression atmospherique.

Les fibres utilisees au cours de cette etude sont des fibres T300 (ex-PAN, orlglne Elf-Aquitaine) brutes de graphitisation, a 6000 filaments unitaires de 7,2 microns de diametre moyen.

2.2. Recherche des conditions de la RCVD

En fonction des resultats thermodynamiques precedents, nous avons considere uniquement les melanges de compositions 2,3 R 5. Dans tous les cas envisages (1300K T 1700K, lmin t 15min.), nous avons mis en evidence la formation de SiC-B.

La morphologie du revetement SiC depend des parametres R, T et t. Les images MEB de la ~ resument les cas extremes rencontres. A T = 1700K et a t = 15min., la fibre devient rigide, les filaments externes du toron sont emprisonnes dans une gangue SiC-B (fig.4a). A cette meme temperature et en reduisant le temps a 3 min., les filaments ne presentent plus de soudure mais sont recouverts par endroit de cristaux de SiC (fig.4b). Cette formation de cristaux s'explique en admettant la reaction:

- (2)

puis celie entre le methane forme et SiC1 4 introduit:

Un revetement homogene et uniforme des fibres a ete obtenu en reduisant la quantite d'hydrogene dans la phase gazeuse a R = 2,3 et en diluant ce melange par de l'argon: ce resultat etait prevu par la thermodynamique puisque cette dilution limite la reaction (2). La

260

morpholgie de la surface des filaments reste alors semblable a celle de la fibre de depart (fig.4c).

2.3. Proprietes des fibres C(SiC)

Les resultats qui sui vent ont ete obtenus sur ces derniE~res

fibres. Elles ont ete recouvertes 3 a ~1 10000C dans un 3mela~¥e caractense 3par ~1= 2,3 (d(H 2 ) = 75cm .min , d(SiCI 4 ) = 32cm .min , 015

340cm .min ). La vitesse de defilement de La fibre est de - -1 m.h ,ce qui correspond a un temps de traitement de 2 min.

2.3.1. Caracterisation du depot

La presence de· SiC-B est mise en evidence par diffraction de rayons X et par spectrometrie Raman. Le recouvrement uniforme des filaments est prouve par MEB 11 partir du residu de combustion du carbone obtenu apres chauffage 11 l'air 11 7000C pendant 7 jours ( i!.s...:..2 ).

Les caracteristiques de 3 fibres C(SiC), referees Fl, F2, F3, sont resumees dans Ie tableau 1. L'epaisseur du revetement ne peut pas etre directement mesuree par MEB, elle a ete evaluee par analyse chimique du silicium et du carbone.

La contrainte 11 rupture et Ie module d'elasticite ont ete determines a partir de tests de rupture en traction ~pr monofilaments de 20 mm, la vitesse d'etirement etant de 0,1 mm.min • Les resultats exploites par la statistique de Weibull mont rent qu'un depot de 0,05 micron n'affecte que tres legerement les caracteristiques de La fibre initiale.

2.3.2. Efficacite du revetement

L'efficacite du revetement SiC a ete demontre par des experiences d'oxydation 11 l'air. Elles ont consiste 11 mesurer la perte de masse de la fibre chauffee 11 l'air, 11 temperature constante, en fonction du temps. A 450oC, la fibre de carbone perd 60% de sa masse en 10 h, la fibre F1 garde pratiquement sa masse. A 600oC, au bout d'une heure la fibre de carbone est consumee 11 85% alors que la masse de la fibre F1 n'est reduite que de quelque %. La ~ permet de comparer les fibres T300, F1, F2 et F3: il apparait que Ie depot de SiC .(F3) n'assure pas la meilleure protection. Ceci est certainement 11 relier 11 la presence de microfissures dans le revetement le plus epais, qui amene des chemins preferentiels de diffusion et une fragilisation de la fibre.

CONCLUSION

Le procede de RCVD permet Ie traitement en continu de surface des fibres de carbone 11 la pression atmospherique. Il est experimente sur la fibre PAN T300. Le carbone de la fibre reagit sur Ie melange SiCl4-H 2 11 une temperature superieure 11 10000C pour conduire en surface de chaque filament unitaire 11 un depot mince, uniforme et adherent de SiC-B. Les proprietes mecaniques des fibres C(SiC) sont

261

VOlslnes de celles de la fibre initiale. 11 est montre que la presence de 0,05 micron de SiC permet de reduire considerablement la vitesse de combustion de la fibre.

Remerciements: Nous tenons a remercier la Direction des Recherches et Etudes Techniques pour l'aide financiere qu'elle nous a accordee pour realiser ce travail.

Bibliographie

1- M.F. AMATEAU, J. Compos. Mater.,10 (1976) 279

2- H. KAHN, Met. Trans., Sep, Vol 7A, n09, (1977) 1928

3- J.C. VIALA, J. BOUIX, Mat. Chern. Phys, 11 (1984) 104

4- R. WARREN, C.H. ANDERSSON, Composites 15 (1984) 101

5- V. LAURENT, These de Doctorat Grenoble (1988)

6- J. BOUIX, M. CROMER, J. DAZORD, H. MOURICHOUX, J.L. PONTHENIER, J.P. SCHARFF, C. VINCENT, H. VINCENT, Rev. Int. Hautes Temper. Refract. Fr., (1987) 5

7 -J.C. VIALA, J. BOUIX, H. VINCENT, C. VINCENT, J.L. PONTHENIER, J.DAZORD , Brevet CNRS n08617157 (4112186)

Fibre T DTEX %Si %C %N %0 e u E m °C nm M~a GPa Weibull

T300 386 91,62 7,78 0,19 3150 200 6,9 Fl 1070 378 1,1 94,0 4,26 0,29 20 2400 196 2,5 F2 1270 384 3,2 94,82 2,25 0,18 50 2600 208 2,2 F3 1470 388 6,7 91,71 2,15 0,30 90 1400 187 3,0

Tab.1: Caracteristiques des fibres

262

102,-------------------------,

T.1700K

"SICI •• ,

10" "HZ· ,

10'· Q

/ /

/ HZ

.C

102,.-------------­

'I,

Q

c · ?'

Fig.1: Variation du nombre de moles d'une espece i pour 1 mole de SiC1 4 avec Q pour T = 1700 K et pour: R = 5 et R = 14,2

+

1100 1300 1500 1700 T/K

Fig.2: Limite du domaine de depot de silicium

' • • " IM 11".4 • • ",.

U.--~--.:;";,.:I--_-...;I.:;.'.:.'-----

... u

1.1

~!!!~ ~_J_IO_'...;~:::::::::::=-:=-=-= ~_~"_.~K ______________ ___ ___ IIIU

.~-------------------------. • 1.1 01 1.1

"

II

~ . . ~ II

I.'

263

,. •• " 1.",, 1" . . ...... 1

I .U I_U • /----;=======::::;::::=;::;:;-;:::---- -<---- 1

~---- - I- I~;,:;;: • ., ~ ~ I: •

_ uru: :

-..-- .,,,~ : -. - '""-..,:

1-. _____ •. 1 ___ _ oj ---- " t

I" '~ " •• '1'1"1'

Fig.3: Evolution des ~aux de conversion de SiC1 4 en SiC (p) et de C en SiC ('I) en fonction de la dilution du melange H2-SiC1 4 par l'argon, a P = 1 atm.

a: 1700K, 15 min b: 1700K, 3 min

c:

Fig.4: Images MEB de la surface de filaments C(SiC)

264

Fig.S: Residu d'oxydation de la fibre Fl(7000C, 7 jours)

Am r-----~--____ --__ ----~----_, m.

100

80

80

40

20

Ox),dalion air

600 · C . lh

Fig.6: Perte de masse de differentes fibres apres traitement a 600°C a 1 '.ai r.

SPUTTER DEPOSITION OF DIFFUSION BARRIER COATINGS ON SiC MONOFILAMENTS FOR USE IN Ti·BASED COMPOSITES

R. KIESCHKE, R. SOMEKH, T.w. CLYNE

University of Cambridge Department of Materials Science and Metallurgy

Pembroke Street - CB2 3QZ CAMBRIDGE - England

ABSTRACT

Procedures are described for the formation of thin layers on monofilaments by sputter deposition, with the aim of controlling high temperature degradation by interfacial chemical reaction. Factors affecting material selection and coating structure are briefly considered. Advantages are identified in having a relatively high fibre temperature during deposition, in encouraging stored stored compressive stresses and in generating a duplex coating structure, comprising an inner layer of metal and an outer layer of oxide.

1. INTRODUCTION

The development of composite materials based on reactive metal matrices has been hindered largely because of progressive chemical reactions between matrix and reinforcement. The matrices of primary interest in this context are Ti, Ni & Mg alloys. Of the currently-available fibrous reinforcements materials, SiC is one of the most stable. Even with this system, however, there are still chemical activity problems. In contact with Ti at elevated temperatUre metal silicide and carbide are formed. It should be emphasized that reaction may occur both during the fabrication of composites and under service conditions. The diffusion bonding process currently employed in fabricating titanium-based materials, for example, involves a thermal history sufficient to produce significant interfacial reaction [2,3].

Fibre coatings have been produced for various purposes, including the encouragement of wetting [4], using a wide range of deposition techniques. Thin Y203 coatings have been used to slow reactant transpon kinetics in the Ni/SiC system [5,6]. These have been shown to inhibit reaction, but are prone to damage during thermal cycling. Coatings of pure metals and materials such as TiC, TiN, B4C &TiB2 have also been used in the AVSiC system to improve wetting behaviour and reduce the chemical attack of a C-rich SiC surface layer[7]. Some coatings can be produced by CVD [8], but problems often arise in finding a suitable carrier gas and establishing a deposition

266

procedure for selected coating materials. In addition, control over factors such as deposition temperature, and the microstructure and stress state of the deposit, will in general be limited. Sputter deposition, on the other hand, is a very versatile coating technique, offering good control over deposit density[9, 10] , stress state[ll], composition and thickness[12].

2. BARRIER REQUIREMENTS

2.1. Material Selection

In the process of barrier selection, criteria for thermodynamic stability in the temperature range of interest are clearly of importance. From this point of view, highly stable oxides emerge as strong contenders. Using the Gibbs free energy of formation of several candidate oxides (Y 203, Hf~, & ZID2) and comparing this with the free energy of formation of Ti02 provides a simple basis for assessment of the stability of these materials for use in a titanium matrix. The relevant thermodynamic data are presented in T bl I F tho Y 0 th bl f h bo 'd a e . rom is, 2 )3 emerges as emoststa eo tea veOXl es. Elemen In oxide of X

X cation ionic radius defects ~Gl()()()f, DGnq~f) D~q~f) charge (pm) (kJ rnol-) Ti +4 68 - -710 - -

Mg +2 66 - -996 - -Li +1 68 - -935 Y +3 89 o defective -1080 1.3 x 10-21 1.0 x 10-16 ZI +4 79 o defective -840 1.8 x 10-26 4.1 x 10-14 Hf +4 78 o defective -892 Not Available Not Available

Table I. Summary of Thermodynamic & Diffusion Data [13·15]

Clearly, the barrier must impair through-thickness transport of reactants across the interface .. Transport rate predictions are hindered by lack of diffusivity data, and simplifying assumptions are necessary. Firstly. given that progressive reaction is not possible without the transport of at least two of the reactants through the barrier, one of these reactants is expected to be cationic in character. It is assumed that this reactant will diffuse on the cation sublattice of the candidate oxide, at a rate comparable to that of the oxide's own cation self-diffusivity. (This is likely to be the rate-determining process, as anion transport is faster in these oxides.) Similarity of cation size and charge may ensure that the reacting cation diffuses no faster than the self-cation, although it may also encourage a high solubility. Table I gives values of the cation size, charge and also the self-cation diffusivities for the candidate oxides. In the candidate oxides, the defect structure is such that the majority of vacancies lie on the anion (oxygen) sublattice, rendering the structure oxygen-defective. From Table I, ZI02 (and probably Hf02, which is very similar in many respects) appear attractive in terms of low diffusivity, but may in some circumstances be insufficiently stable.

On balance Y 203 appears the most promising choice of barrier from the above oxides. In an attempt to quantify the diffusion characteristics of interest, a finite differel'ce model was used to calculate the total flux of reactant (in this case Ti) transported through the barrier in 1000 hrs at temperatures from 1000K to 1200K. This flux was then converted to a reaction product thickness via a mean reaction product stoichiomeuy. Figure 1 shows the reaction product thickness as a function of barrier thickness. A ,falue must be assumed for the maximum solubility of Ti in Y 203: as this appears to be unavailable, a rather pessimistic (high) value of 10 atom% was employed in figure 1. Also shown for the 1200 K curve is the effect of reducing this value to 5%. It

267

should be emphasized that the the predicted rates are lower bounds, as they use self-cation diffusivity values and take no account of possible short-circuit diffusion paths. The plots in figure 1 allow estimates of a minimum barrier thickness, given that, on the basis of critical flaw size concepts[16,17], a target maximum thickness of reaction product (in which cracks are likely to appear) is probably of the order of 300-500 nm. This is clearly rather open-ended, given the various uncertainties, but a barrier layer of at least 500 nm in thickness would appear to be advisable.

Sputter deposition is attractive in allowing a degree of control over the deposit stress state by adjustment of the pressure of the sputtering gas [11]. If this is low, neutralized backscattered argon atoms suffer fewer scattering events before hitting the substrate, so that the average incident impact velocity is increased. This substrate bombardment tends to generate compressive stresses in the deposit ("atomic-peening"). These compressive stresses are superposed on the natural tensile stresses present in a growing undercooled fUm (due to excess vacancy incorporation). The as-deposited stress will therefore change from tensile at high sputtering gas pressures to compressive at low gas pressures.

2.2. Duplex Barriers

Mechanical and other considerations have led us to consider the merits of a duplex barrier [18], consisting of a metallic layer adjacent to the fibre with an overlayer of the metal oxide. This ductile metallic underlayer offers the possibility of reducing the deleterious effect of brittle interfacial layers, by acting to prevent crack propagation into the fibre. In the case of titanium, since a residual oxygen content is always present in the matrix, a suitable choice of metaVoxide system may also offer the prospect of oxide self­repair by means of oxygen gettering from the matrix. If a Y!Y203 layer is chosen, thermodynamics predicts that Y in contact with a titanium/oxygen solution should getter oxygen from the titanium matrix, leading to the formation of Y 203 on the Y rri interface. The viability of forming Y 203 in Ti/Y /0 solutions has been confirmed by Y 203 dispersion strengthening work [19].

The sputter deposition process offers the possibility of overlaying an oxide via a reactive sputtering technique [12]. Hence, after depositing an initial layer of Y, a covering layer of Y 203 could be applied (without removing the fibre from the deposition chamber) during a second sputtering run. Partial layer oxidation also appears attractive. However, given that the (compressive) stresses involved in oxidation may be at least partially relaxed at the heat treatment temperature, and that CTE values for SiC, Y 203 & Y are 4, 8 & 11 Il(K-I) respectively, there would be concern that tensile stresses generated on cooling from the heat treatment temperature may lead to hoop and axial cracking. It may therefore be preferable to encourage in-situ formation of a layer of Y 203. Since the diffusivity of oxygen in Y203 is about 105 times greater than that of Y (and presumably other cations) in Y203 [15], the rate of oxide growth is expected to be greater than the rate of Ti penetration, so that the diffusing titanium would always "see" a Y 203 layer. Typical oxygen contents in Ti are of the order of 2000 ppm [20]. With this oxygen content, and a 30% volume fraction of fibres, the Y layer would become oxidised to a depth of 500 nm if all this oxygen were removed from the matrix. Furthermore, using a value for the diffusivity of oxygen in n-Ti at 1000 K of about 1.7 x 10.13 m2 s·I, the necessary diffusion through the matrix would take place in an hour while diffusion through the growing Y 203 layer would also need about the same time. Therefore the transport of oxygen through the matrix will not limit the oxidation rate. Sastry [19] has shown that the presence of Y can reduce the concentration of oxygen in titanium, down to 100 ppm at 1000 K. (Below this level the entropy reduction is presumably too great to allow further gettering).

268

3. EXPERIMENTAL

3.1. Fibre Coating

Sputter deposition was carried out with the sets of apparatus shown in figure 2. With the batch rig, the fibre-magnetron distance had to be relatively large, limiting

the heating effect from the energetic magnetron discharge (so that the fibres were close to ambient temperature during deposition). Deposits formed on such cold substrates are expected to exhibit low densities because of slow deposit bulk and surface diffusivities. Furthermore, beneficial atomic argon bombardment of the deposit is expected to be limited at such large stand-off distances. Using the continuous rig, on the other hand, the target/substrate distance was small enough to allow considerable fibre heating from the magnetron discharges and from general bombardment Estimated fibre temperatures were of the order of 500'C in the deposition zone. The two reels were mechanically linked, providing a smoother motion and allowing the fibre to be transported in both directions. Typically, the fibre feed velocity would be -0.3 - 3 mm s-l.

3.2. Sputter Deposition

A single set of conditions was employed with the batch rig. The continuous rig was used to provide coated fibre with varying deposition conditions on short sections along its length. The fibre used was a 100 ~m diameter, tungsten-cored SiC monofilament ("Sigma" fibre). Sputter deposition was carried out using an yttrium target material. Various deposition conditions have been employed.

3.3. Generation of Duplex layers

Several lengths of Y -coated fibre were evacuated to a pressure of less than 10-5 torr and sealed in a SiCh ampoule with some Ti powder. This tube was then heat treated at 1173 K for 1 hour. The residual oxygen content was sufficient to cause partial oxidation of the coating. In addition, duplex layers were produced by reactive sputter deposition of Y 203 overlayers, which takes place in two stages. Firstly, the surface of the Y target material oxidizes, and sputtering occurs from a Y 203 target. Secondly, oxygen is available in the chamber to encourage the formation of stoichiometric Y 203 on the substrate if the sputtering yields of Y and 0 differ greatly.

3.4. Examination of Fibre Coatings

The fibre coatings were examined (on fractured fibres) in "Camscan" series 4 and series 2 SEMs, giving information on coating thickness, integrity, topography and adhesion. TEM examination was carried out using Philips 400T and JEOL 120CX microscopes. TEM specimen preparation followed a route of Ni plating a fibre array to give a total thickness of 3 mm, followed by a process of slicing, grinding to about 60 ~m thickness, dimpling, and ion milling to perforation using 6 ke V argon ions at an incident angle of 15".

4. RESULTS AND DISCUSSION

4.1. Effect of Deposition Variables

The coatings produced on the batch rig exhibited some evidence of inhomogeneities and poor adhesion; for example, a spalled region can be seen in Fig. 3. Furthermore, TEM micrographs, such as that shown in Fig. 4., revealed a very fine grain size and

269

extensive fine scale porosity. Fibres produced on the continuous rig all appear to show improved adhesion over the batch process coatings. For example, Fig. S illustrates portions of coating adhering to a fragmented region of the fibre. This is consistent with the fibre/coating adhesion being improved at higher substrate (fibre) temperatures, assuming that this will increase with increasing total magnetron power. This is expected because the fibre heating mechanisms (electronic and atomic bombardment, and the heat of dissociation) will be expected to increase with increasing power, while axial conductive heat flow along the fibre length will remain small. An increased fibre temperature is clearly desirable in terms of providing large enough bulk and surface diffusion rates to allow dense deposits to be formed. In addition, a relatively large grain size is probably desirable for efficient functioning of the deposit as a diffusion barrier. Work: is currently in hand to examine continuous coatings using the TEM.

4.2. Duplex Coating Production

Fig. 6 illustrates a duplex structure obtained by partial oxidation. Problems arise, however, from differential thermal contraction stresses on cooling from the heat treatment temperature. Fig. 7 shows the axial and hoop cracking that occurs on cooling from the heat treatment temperature (1173 K).

An indication of the magnitude of the stresses giving rise to this effect can be obtained by employing an analytical model developed for continuous coaxial cylinders [21]. In Fig. 8, the effect of a temperature decrease of 500 K is shown for a coated fibre with and without the matrix present. In the latter case the predicted hoop and axial stresses in the coating would be sufficient to cause the observed cracking. The fact that the coatings did not crack on cooling from the deposition temperature with the continuous rig (-SOO'C) suggests the introduction of high compressive stresses during deposition and this is also consistent with x-ray measurements [21].

S. CONCLUSIONS

A rationale has been presented for optimisation of the characteristics of barrier coatings on SiC monofilaments, designed to protect the fibre against chemical attack in reactive matrices such as titanium. The following points have been identified:

(a) Thermodynamic stability considerations severely limit the choice of material. yttria has been identified as one of the few compounds likely to be sufficiently unreactive.

(b) In order adequately to impair transport of reacting species, the thickness of the layer probably needs to be at least 500 nm.

(c) It is proposed that a duplex barrier, for example one with an inner layer of metallic Y and an outer layer of Y z03, may offer certain advantages. These include decreased danger of crack propagation from oxide layer to fibre and potential, particularly in Ti matrices, for "self-healing" of damage to the Y 20:3 layer as a result of the gettering of dissolved oxygen from the matrix by the exposed Y. It is confirmed that the rate of transport of 0 through a Ti matrix will not limit the operation of this mechanism.

(d) Coatings have been produced by sputter deposition, a process with scope for control over the microstructure and stress state of deposited material. Duplex coatings have been produced by reactive sputtering of Y 20:3 onto a Y underlayer and by partial oxidation of a sputtered Y layer. The latter can be accomplished both by heating the coated fibre in oxygen and by an in-situ heat treatment after composite fabrication in the presence of dissolved oxygen in the matrix.

270

(e) A higher fibre temperature during sputtering has been identified as generating improvements in the microstructural integrity of the coating and its adhesion to the fibre surface.

(f) Stress state calculations suggest that high compressive stored stress induced in the coating during sputtering is necessary in order to avoid cracking during post­deposition cooling to room temperature.

ACKNOWLEDGEMENTS

Financial support from Rolls Royce pIc for one of us (RRK) is gratefully acknowledged. In addition, the authors are grateful to P. Doorbar and 1. Hooker for useful discussions.

REFERENCES

1. P. Marteneau, M. Lahaye, R. Pailler, R. Naslain, M. Couzi, F. Cruege, I. Mater. Sci. 12.(1984) pp2731-2748.

2. C.G. Rhodes & R.A. Spurling in "Recent Advances in Composites in the United States and Japan", ASTM SIP 684, J.R. Vison & M. Taya (eds.), Am. Soc. Testing & Mats. Philadelphia, (1985),585-599.

3. P.R. Smith & F.H. Froes, J. Met. March 1984 ppI9-25. 4. H.A. Katzman, J. Mat. Sci. 22(1987) 144-148. 5. R.L. Mehan, M.R. Jackson & M.D. McConnell, J. Mater. Sci. 18(1983) pp3195-3205. 6. R.L. Mehan & M.R. Jackson Ceram. Eng. Sci. Proc.l(1982) pp484-503. 7. F.E. Warner & S.R. Nun, Ceram. Eng. Sci. Proc.l(1980) pp709-719. 8. A.G. Evans, Proc. 9th Risl'llnt. Symp. on "Mechanical and Physical Behaviour of Metallic and

Ceramic Composites" 13-34. 9. J.A. Thornton, J. Vac. Sci. 2.1(3) 833, 1982. 10. I.A. Thornton, J. Vac. Sci. 12.(2) 205, 1981. 11. D.W. Hoffman, Proc. 7th ICVM, Iron & Steel Inst. Japan, 1982, Tokyo, Japan. 12. Z.H. Barber, Ph.D. Thesis, University of Cambridge, Cambridge, UK. 13. M.P. Berard & D.R. Wilder, J. Appl. Phys. 31(1963),2318. 14. M.F. Berard & D.R. Wilder" J. Amer. Cerarn. Soc. 52.(1969), 85. 15. P. Kofstad, "Nonstoichiometry, Diffusion and Electrical Conductivity in Binary Metal Oxides",

Wiley Interscience, 1972, USA, 268-274. 16. S. Ochiai & Y. Marakami, I. Mater. Sci. 12.(1979) pp831-840. 17. S. Ochiai, S. Urakawa, K. Ameyama, Y. Marakami, Met. Trans . .llA(1980) pp525-530. 18. UK Patent No. 8812556.2 "Improvements in or Relating to Coated Fibres for Use in a Metal

Matrix and in a Composite Structure", 26th May 1988. 19. S.M.L. Sastry, PJ. Meschter, J.E. ONeal, Met. trans. A. 1M ppI451-1463. 20. 1.1. Polmear, in "Light alloys - Metallurgy of the light metals" ch 6, p162, Arnold 1986. 21. C.M. Warwick, R.R. Kieschke & T.W. Clyne, to be submitted to Met. Trans. A.

1 0.8 ill

~ ~ 0.6

~ '8 0.4 II: c .2 i 0.2 ;;.

Fig. 1. Reaction Product Thickness After 1000 hours as a Function of

Y:z03 Barrirz Thickness for a Titanium Reactant. (Maximum

concenttation 10%)

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4

Barrier Thickness (J.lm)

<a) Balch

IOJ.Im (a)

(b) Continuous Fig. 2. Schematic diagrams of the Coating Rigs

3J.1m (b)

Fig. 3. SEM micrographs of the batch Coating

271

272

500nm

111m

Fig. 4. 1EM micrograph of the batch coating.

Fig. 5. SEM micrograph of continuous coating

1OJ.lffi Fig. 6. SEM micrograph of the partially oxidised duplex

coating. Fig. 7. SEM micrograph of partially oxidised

coating. showing axial and hoop cracking.

o (Mr.)

"'" ... ,.. '"

·100

."" .,.. -.JaI -.". ..,.

w SIC

I ; I: o.~

, w,mmmnomn .. p, W

;~ 1°, .... ; ---o=-.--~n; ""- ... -( dT ;;; -SOD K

Fig. 8. Stress Slate Predictions (referring to the effect of a 500K temperature decrease) for a Coated Fibre Both With and Without Surrounding Matrix

ABSTRACT

SIMS ANALYSIS OF SiC COATED AND UNCOATED NICALON FIBERS

M. LANCIN. J. S. BOUR

CNRS - Laboratoire de Physique des Materiaux 1 Place A. Briand - 92190 MEUDON - France

Pyrolytic SiC deposits were performed on Nicalon 202 SiC fibers using either a static method designed for fiber structure infiltration or a dynamical method designed for fiber coating. A Carbon deposit was sometimes introduced between the fiber and the SiC matrix. The coated fibers were analysed by Secondary Ion Mass Spectrometry (SIMS). The depth profiles obtained show that (i) the fibers retain their composition after the process (static or dynamical), (ii) the deposits prepared using the static method exhibit an heterogeneity of microstructure and composition; these changes are due to a transitory stage at the beginning of the CVD reaction and (iii) the deposits prepared using the dynamical method exhibit a constant carbon content higher than the stoechiometric one (C/Si = 1.4).

INTRODUCTION

ONERA Chatillon has developped Chemical Vapor Deposition Process (CVD) for the infiltration of fiber structures ("static" method /1/) and for the coating of yarns ("dynamical" method). The aim of these processings is to prepare materials for thermomechanical applications. Among the components which can be prepared by CVD, SiC is one of the most interesting due to the high potentiality of both SiC/SiC composites and SiC coatings. Pyrolytic SiC deposits were thus realized by ONERA on substrates which were either fixed in the reaction chamber (static method:a) or moved through the reaction chamber (dynamical method:b).

The chemical composition of the deposits and the interdiffusion fiber/matrix were studied on Nicalon 202 fibers coated with SiC. In order to separate the contributions of the heat treatment during CVO and of the interdiffusion fiber/matrix we studied the chemical composition of :

274

(i) the as received Nicalon 202 fibers, (ii) the fibers annealed with the condition of the CVD heat treatment, (iii) the fibers coated with SiC-CVD.

EXPERIMENTAL PROCEDURE

The samples are segments of fibers which were coated using either method a or b.

The chemical composition was determined by ion sputter depth profiling in a Cameca IMS3F. The lateral resolution was equal to 1.4 ± 0.5 ~ ,the resolution in depth was equal to 3 nm and the sensitivity was superior to 10- 4 at% /1/. An experimental procedure was developped (i) to realize radial analyses and (ii)to limit the sputtering during the alignment of the ionic optics before the analysis and therefore start the me"asurements at a depth of a few nm /3/. Standards were used to determined the chemical composition of the materials /2/.

RESULTS

Nicalon 202 fibers are rich in carbon and oxygen as it was previously found while using Scanning Auger Microprobe (SAM) /4/ and electron probe /5/ technics. Extensive SIMS analyses of the fibers have shown that the chemical composition is constant along their diameter (Fig.l). On the contrary, the composition changes along their axes as shown by the scattering of the following ratios : C/Si = 1.3 ± 0.7 and O/S; = 0.45 ± 0.15. Assuming that the fibers are made of SiC , Si02 and C the mean molar composition is : SiC = 51%, C = 34% and S;02 = 15% /3/.

The heat treatments equivalent to those sustained by the fibers during the CVD processes (a or b) do not change significantly the fiber composition, at least in the bulk as the very fiber surface composition was not investigated.

~

lSI lSI -x

10 3

.... 10 2

<t-O

'" H , H

Masse

~---

x.

ref: 28

C-/S.-

0-/ 5 ;-

1- z tm.

Figure 1: SIMS analysis of a Nicalon 202 fiber. Variation of the ratio of the ionic intensities versus depth.

275

Coated fibers prepared using the static method were extensively studied.

Two types of coated fibers were analysed :Nicalon 202 fibers/C-CVO/SiC-CVO and Nicalon 202 fibers/SiC-CVO. The C-CVO deposits were performed at T = 1373K under P = 133Pa; their thickness ranged from 65 to 90 nm. The SiC-CVO deposits were realized at T = 1523K under P = 66 Pa; their thickness ranged from 1 to 8 ~.

All the materials were characterized by SIMS, Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM) /2,3,6/. Electron Energy Loss Spectroscopy analyses (EELS) were also performed on some SiC/C/SiC coated fibers /7/. In both SiC/C/SiC and SiC/SiC composites, characteristic features can be pointed out (fig. 2 & 3) : (i) changes in fiber composition were detected by SIMS next to the interface: particularly a decrease of the oxygen content occurs over

N I~; we can therefore conclude that oxygen diffuses out of the fibers during the process. No correlated changes in the microstructure and the phase composition were detected by MET and EELS as it was observed by Rawlin et al. /8/. (ii) the CVO deposits always consist of two layers so called CVOI and CV02 (formed during the early stage of the CVO). These layers exhibit distinct microstructure and composition.

The CV02 layer composition is reproducible and slightly sub-stoechiometric in carbon.

The thickness of the CVOI layer and its composition are not reproducible. The carbon and oxygen contents are higher in the CVOI than in the CV02. The changes in the SiC-CVO composition near the interface cannot be due only to the interdiffusion between fibers and the SiC deposits.This layer is probably formed during a transitory stage which occurs at the beginning of the CVD reaction.

• • I. I. b)

0.) C C- •

I. s

Si-I.

~.

~ •• 4 iii.e I ! CD

!' .

· .. · • Mel' ~ fjber · • _I. ~I • .!i

I. • c.V:Pi. -I C"JH.

\J- I.

s elf:P f-,. -. ell]) z.. • 30mn t I •

3.um • Figure 2: Depth profile of a coated fiber SiC/C/SiC obtained by SIMS. Variation of the ionic intensities versus time and of the ionic intensity ratios versus depth. The thickness of the carbon deposit was equal to 65 nm.

276

CNRS-MEUDON DEPTH PRO~ILE IB3~~I~M~S~3~~ ______________________________________ ~

c-/sz ~ IB2~kl __ ~~~~ __________ ~~-t----------~======:] lSI lSI ... X

" IB 1 <t­O I..

H

" B HIB

IB- 1

f~bre

CVD2

Figure 3 : Depth profile of a coated fiber SiC/SiC obtained by SIMS. Variation of the ionic intensity ratios versus depth.

The change in the fiber surface composition as well as the lack of reproducibility of the CVD deposits are detrimental to the realization of composites exhiblting well defined characteristics. Therefore, the CVD parameters were more strictly controlled and new coated fiber segments prepared. SIMS analyses showed that (fig.4) : (i) the fibers retain their original composition after the process, (ii) changes in the carbon and oxygen contents of the deposits are still observed in the early stage of the CVD process. This demonstrates that the changes in the SiC composition are not due to the interdiffusion fiber/deposit but originate in a transitory stage at the beginning of the reaction which cannot be avoided.

Coated fibers prepared using the dynamical method. The coatings (e N 5 ~) were realized at T = 1523K under P = 66Pa while fibers segments were moved through the reaction chamber. SIMS analyses (fig.5) demonstrate that : (i) the process does not change the fiber composition, (ii) in the SiC deposit the carbon content is higher than the stoechiometric one (C/Si = 1.4). It does not change during the process whereas the oxygen content still varies.

277

1a 2 j ,ell]):!.1 C - / s-: j ~= ~ IS) / 0-/S-: IS)

x ~ 1

1a ...

1 11 '-

H , H

100 ~ J~

1 L---C"J)2.

-1· 0,5 ~,o "", S 2,0 1,5 .3,0 3,5 4,0 I I I I I I I 10 __ . ,m.

Figure 4: Depth profile of a coated fiber SiC/SiC obtained by SIMS. Variation of the ionic intensity ratios versus depth.

,

j J

fm l

Figure 5 : Depth profile of a coated fiber SiC/SiC obtained by SIMS. Variation of the ionic intensity ratio versus depth.

278

CONCLUSION

SIMS analyses were realized on Nicalon 202 fibers coated by SiC or by C and SiC using either a static or a dynamical method. When the static method is used, a transitory stage occurs at the beginning of the CVD reaction. It results in an heterogeneity both of the composition and of the microstructure of the SiC deposit. With the dynamical method such a transitory stage seems to be avoided. Therefore, this last method appears to be promising. In both cases, no change in the fiber composition and structure was observed after the process. Therefore, the fiber should retain its properties. However to verify this assertion, new investigations of the interfaces between fiber and deposit by HREM correlated to mechanical tests are necessary.

ACKNOWLEDGEMENTS

J.S. BOUR wish to thank the DRET for a maintenance grant and the financial support. Authors are debtful to M. SCHUHMACHER who gave them useful advices for SIMS analyses and to L. GRATEAU and M. PARLIER for providing the materials.

REFERENCES

1 - Parlier M., These Ecole Sup. des Mines, Paris (1984). 2 - Anxionnaz F., These Universite Paris VI (1987). 3 - Bour J.S., These Universite Paris XI (1988). 4 - Brennan J.J., Proc. of the Conference on Tailoring Multiphase

and Composite Ceramics. Penn State University (17-19 July 1985). 5 - Lesniewsky, Rapport de Contrat DRET, n084-447 (1986). 6 - Lancin M., Anxionnaz F., Schuhmacher M., Dugne 0., Trebbia

P., Mat. Res. Soc. Symp. Proc. 78 (1987) 231-238. 7 - Anxionnaz F., Schuhmacher M., Trebbia P., Lancin M., J.

Microsc. Spectr. Electron. 11 (1986) 421-427. 8 - Rawlins H.H., Nolan T.A., Stinton D.P., Lowden R.A., Mat.

Res. Soc. Symp. Proc. (1987) 293-230.

THE EFFECT OF SURFACE TREATMENT ON THE INTERFACIAL STRENGTH OF CORROSION RESISTANT GLASS FIBRES IN A

VINYLESTER RESIN

F. JONES, 0 PAWSON

University of Sheffield School of Materials, Northumberland Road, S10 2TZ SHEFFIELD - England

ABSTRACT

The effect of aqueous condit ioning on the bond strength of AR glass fibre-vinyl ester resin composites has been investigated using the embedded single filament tensile test. Two fibre coatings were compared, one containing Al100 (~ amino propyl triethoxysilane) silane coupling agent and the other without. The results provide evidence for the existence of physisorbed and chemisorbed layers on the fibre surface. The former was removed by immersion in warm water whereas the latter is tenaciously bound to the surface, providing protection against boiling water and sulphuric acid.

INTRODUCTION

Silane coupling agents are an important component of modern reinforced plastics and composites. Glass fibres are usually coated with a size immediately after forming, which usually consists of a lubricant, a binder and a silane coupling agent. The coupling agent improves the bond between the glass and the resin especially in wet environments. There are several theories which aoH empt to explain the mechanism by which they improve the mechanical properties of glass fibre composites /1/. These include the preferential adsorption, the interpenetrating network, surface wettability, chemical bonding, deformable layer and restrained layer theories.

It is known that coupling agents adsorb onto glass fibres as multilayers. There is evidence to suggest that three regions of differing structures exist /2/. Several techniques including radioactive labelling 12/ and Fourier Transform Infra-red Spectroscopy (FTIR) /3/ have been used to study the structure of the interfacial

280

region. It is generally considered that the coupling agent structure comprises a chemisorbed layer adjacent to the glass surface consisting of a strongly held polysiloxane network with an outer physisorbed layer consisting of about 98% of the coupling agent made up of hydrolysed or partially hydrolysed silane and smaller polysiloxane moledcules. The outer lower molecular weight materials can be easily removed by organic solvents or even prolonged contact with water. In contrast the inner chemisorbed layer can be separated into two components; the outer which can withstand 3 to 4 hours in bOiling water, and a tenaciously held monolayer equivalent which remains chemically bound to the surface after hydrolysis.

The physisorbed layers of silane dissolve into the resin during cure, forming a complex interfacial region of resin silane oligomer immediately around the fibre surface. This region has mechanical properties which differ from the bulk resin and modify the mechanical performance of the composite.

, The aim of this research was to determine the effect of surface treatment and different environments on the interfacial shear strength of AR glass fibres, vinyl ester resin composites using the embedded single filament technique. Fibre glass coated with different silane coupling agents was given different treatments in wet and dry environments prior to immersion in the activated resin. This approach differs from others in that the matrix is constant, any change in interfacial strength is caused directly by a change in the structure of the coupling agent on the glass fibre surface and how it interacts wi th the resin.

I - EXPERIMENTAL

1. 1. Materials

The vinyl ester resin, Derakane 411-45 (Dow Chemical Limited) was cured using 2 phr (part s per hundred) of a 50% methyl ethyl ketone peroxide solut ion (cat alyst M), 1 phr of a cobalt naphthenate solution (Accelerator E), (Scott Bader Limited) and 0.5 phr of dimethyl aniline (10% in styrene). AR glass fibres were available with and without AllOO amino propyl triethoxy silane (Dow Corning Limited) incorporated into a vinyl ester compatible finish.

1.2. Interfacial shear strength measurement

The embedded monofilament tensile test has been utilized to determine the cumulative distribution of critical lengths for a given polymer resin system from which an estimate of the interfacial shear strength was obtained. The treatments given to the fibres prior to testing are shown in Table 1. Fabrication of the specimens and alignment of the fibres was accomplished with the aid of silicone treated steel moulds, 200 rom by 10 mm by 2 rom. The moulds were filled to half the final level with the activated resin using a glass dipstick to direct the resin to all parts of the cavity. Single filaments were selected at random, care was taken not to touch the central part of the fi bres. The filament s were mounted in the centre

281

ot tne mOUld; subsequently tne remainder of the activated resin was then poured into the mould. This method ensures that the fibre lies in the centre of the test piece, parallel to the edges and does not come into contact with the silicone release agent. Specimens prepared in this way behaved in the same way as those in which all the resin was added simultaneously. The test pieces were cured at room temperature for 24 hrs, t hen cut int 0 tensile specimens of dimensions 65 DUD by 10 mm by 2 DUD using a diamond impregnated rotary wheel. To reduce the risk of grip initiated premature failure, the thickness of the central portion was waisted down to 1 DUD and the outer portion to 1. 5 DIll. A Cerium oxide polishing wheel was then used to produce a transparent finished product. The test pieces were pulled in uniaxial tension on a Mayes Universal testing machine at 2 mm/min until the number of fractures within the filament became constant. This was confirmed using a Carl Zeiss optical microscope attached to the testing rig.

The fragment lengths werre measured individually with a travelling microscope. The critical transfer length L. is given by

Lc = 4L13

where L is the mean fragment length. The interfacial shear strength, ~, is calculated according to the equation

't = a,u d/2L

where a,u is the fibre tensile strength at the critical transfer length and d is the fibre diameter. The tensile strength at the gauge length was estimated from the single filament strengths at gauge lengths of 6.35 DUD using Wei bull statistics.

1. 3. Measurement of single fibre tensile strength

Fibre glass tows were given the treatments described 1n the table. Monofilaments were selected at random from the glass tows and mounted on window cards using card endtabs to secure them in place at a gauge length of 6.35 DUD. The fibrediameters were then measured using the laser scattering technique /4/. Tests were performed on an Instron TT-CM-L testing machine using a crosshead speed of 5 mmlmin and a load cell with 100 g full scale deflection. The sets of data at this gauge length conformed well to a two parameter Weibull distribution /6/. Weibull statistics were used to predict the tensile strength at the critical length.

RESULTS AND DISCUSSION

2. 1. AR glass coated with AIIOO coupling agent

One of the primary functions of a silane coupling agent is to improve the retention of the mechanical properties of a composite in wet conditions. To explore the role of the silane coupling agent, AR glass fibres coated with a size, with (C) and without (NC) the inclusion of the AIIOO coupling agent were subjected to the series of treatments shown in the table. The presence of the coupling agent

282

protects the fibres in all the aqueous environments. It appears that AIIOO in combination either with or without the sizing forms a tenacious layer which is not removed by boiling water. Under the same conditions, AR glass coated with A172 coupling agent (a vinyl functional coupling agent) performed better than the NC fibres but not as well as the C fibres 16/. It is well known that the coupling agent probably forms chemical bonds with the f.ibre surface through Si-O-Si linkages. Thre chemisorbed layer of the AIIOO does offer more protection than the A172 16/, a possible explanation for this could be the formation of hydrogen bonds between the amino functional group of the Al100 and the glass surface. Lee /71 suggested that because the critical surface tension of a f11111 of AIIOO on a glass surface was low, it must be adsorbed with the unpaired electrons on the N atom orientated towards or parallel to the surface. Furthermore Koenig et al. 131 used Fourier Transform Infrared Spectroscopy to show that on a silica surface the NH deformation band near 1561 cm-' indicated a strong interaction with the surface silanol groups; AR glass contains over 70% silica.

The low value of t for the as-received C fibres 1s probably due to the presence of a physisorbed layer of coupling agent. It has been reported that the physisorbed layer can reduce the mechanical properties of a composite 151. The liquid resin can easily penetrate the physisorbed layer and copolymerize with functional groups of the coupling agent, but cannot penetrate the chemisorbed layer. Thus an interpenetrating network is formed between the resin and the coupling agent; this interfacial region will have inferior properties to the bulk resin. Removal of the"physisorbed layer does increase the t, thus it appears that the weakest part of the interface is probably the interface between the physisorbed and the chemisorbed region. After Treatment I, the t is reduced to the level of the NC fibres (since drying the fibres increases the '1:, it is likely that this treatment has removed the physisorbed layer). These specimens were prepared from air-dried fibres and the test pieces were tested within 36 hrs of fabrication, thus any water would not have had time to diffuse away from the fibre surface. Although the coupling agent does protect the fibre from moisture, there is a slight drop in the strength of the fibre indicating that some water molecules are able to break through the chemisorbed layer to the glass surface. When coupling agents initially bond to the surface, the hydrated silanol bonds to metal oxides on the glass surface through a reversible reaction with the elimination of water. Reaction with the surface may be through hydrogen or siloxane bonding depending on the severity of the drying conditions. Each bond formed between the coupling agent and the glass surface is hydrolyzable; it is possible that an excess of water at the interface could upset the equilibrium, and under the shear stress conditions of the single fibre test there would be a shift towards free hydroxyl groups and a subsequent reduction in the interfacial shear strength. Following drying in Treatment 2, the large increase in interfacial strength can be understood by the reformation of the original or new bonds with free silanol groups. This adapt ion of Plueddemann's reversible, hydrolyzable theory of adhesion to hydrophilic surfaces Ill, helps explain the slight drop and subsequent rise in '1:. However the large increase in <p is probably due to the

283

tact that the Physisorbed layer nas been removed by the warm water. To produce the increase in interfacial strength following the removal of the physisorbed layer, the outer portion of the chemisorbed layer must have some of the functional amine groups orientated towards tne resin molecules. Thus it is possible that while some of the coupling agent molecules are bonded to the glass surface through both siloKane linkages and through hydrogen bonding of the amine group with metal oxides on the glass surface; others are only bonded through the former, with the functional groups orientated towards the resin molecule. The coupling agent 1I01ecules attached to the glass surface through siloxane linkages and hydrogen bonding could be connected to other coupling agent molecules through free silanol groups. These would also have their functional amine gorups orientated towards the resin molecules.

Treatment 3 did not reduce afu of the fibres any further, however the 't was slightly reduced. This could be due to a partial removal of the outer layer of coupling agent molecules at tached to the surface silanes through the free silanols by hydrolysis.

The presence of a coupling agent on the surface of AR glass increases the stress corrosion resistance significantly in sulphuric acid /6/. To determine the effect of the acid on the interfacial shear strength, the fibres were given Treatment 4. The decrease in strength of both kinds of fibres was similar to the water treatments, however the 't was significantly lower. This is in contrast to the behaviour of unidirectional laminates; the interlaminar shear strength of AR glass/vinyl ester resin specimens immersed in 2M sulphuric acid at 323K for 2 months show a similar reduction in shear strength to those imaersed in deionized water under the same conditions /6/. Thus the reduction in 't is probably due to the degradation of the outer layer of the coupling agent coating which makes contact with the resin. It is also possible that there is a chemical reaction between the functional groups and the acid, which would reduce their ability to react with the resin. It is not clear froll the results whether or not the acid removes any of the physisorbed layer of coupling agent from the glass surface. Silane coupling agents do not function by preventing molecular water from reaching the mineral polymer interface, but by competing with the water molecules for the mineral surface so that water cannot cluster into films or droplets. This suggests·that loss of adhesion always precedes corrosion /1/. This is backed up by the fact that the change in interfacial shear strength in sulphuric acid of the thre kinds of AR glass mirrors their behaviour under stress corrosion conditions /6/.

2.2. AR glass coated with a size in the absence of the silane coupling agent

As expected the presence of a silane coupling agent enhanced the performance of the tensile and interfacial shear strength of the AR glass vinyl ester system under most condit ions. The 't of the as received NC - fibres was only slightly lower than that of the AllOO fibres. The interfacial shear strength could be attributed to a number of differing factors Ill.

284

However the only requirement for bonding a resin to a hydrophilic mineral surface are that it contains functional groups capable of competing with the weak boundary layers such as water 11/. Furthermore it is probable that the silanol equilibrium is less favourable for bond retention in wet conditions and a water resistant bond would only be obtained if the functional group concentration at the interface was sufficient. That Treatment 1 reduces 't, while in contrast to the silane treated fibres Treatment 2 did not increase it implies the existence of chemical bonding within the fibre/silane/resin interface. Treatments 3 and 4 further reduced both parameters; the fibres immersed in acid were slightly stronger than those immersed in boiling water, even though the 't was considerably lower. It is possible that the acid has smoothed the fibre surface, removing stress concentrators, which could explain the slightly higher tensile strength. This would enable the fibre to easily debond from the resin and to slide through the matrix more easily thus reducing 'to Indeed the gradual decrease in 't with increasing severity of hydrolytic treatment could be due to the glass surface becoming increasingly smooth.

3. CONCLUSIONS

The embedded single filament tensile test has been used to measure the interfacial shear strength of A.R. glass fibres in a vinyl ester resin. The coupling agent protected the surface of the fibre from damage and improved the retention of the tensile and interfacial shear strength in wet conditions. The results provide evidence for the existence of physisorbed and chemisorbed layers and the ready removal of the former by immersion in warm water. The latter are tenaciously bonded to the fibre surface, providing protection against boiling water and 2M sulphuric acid.

ACKNOWLEDGEMENTS

We thank Pilkington Reinforcements Ltd and SERC for financial support.

REFERENCES

1. L.J. Broutman and R.H. Krock, Composite Materials, Vol 6 (E.P. Plueddemann ed.) (1974) Academic Press, New York and London. 2. ~E. Schrader and A. Block, J. Polym. Sci. Part C 34 (1971) 281. 3. C.H. Chiang, H. Ishida and J.L. Koenig, J. Colloid Int. Sci. Ii (1980) 396. 4. P.A. Sheard, Ph.D. Thesis, University of Surrey (1987). 5. R.T. Graf, J.L. Koenig and H. Ishida, J. Adhesion ~ (1983) 97. 6. F.R. Jones and D. Pawson, Unpublished results. 7. L. H. Lee, J. Colloid. Int. Sci. 27 (1968) 751.

285

Table

Fibre treatment A1100 coupled fibres (C) Non coupled fibres (NO

o<u (OPa) 't (MPa) CI.,,(OPa) t(MPa)

As received 22.9 14.6 18. 3 14.5

Treatment 1 2 weeks in 19.6 14.2 16. 7 14.3 Deionized water at 323 K

Treatment 2 As above then 16. 7 20.3 14.9 13.7 vacuum dried at 323 K

Treatment 3 2 hours in 17.4 18.9 13.5 13. 1 Deionized water at 373 K, then desiccator dried

Treatment 4 4 weeks in 2 M sulphuric acid 17.2 12.8 14.7 6. 2 lit -298 K, then washed in deionized water

EFFECTS OF MATRIX MICROSTRUCTURE CHANGES AFTER ANNEALING ON FRACTURE PROPERTIES OF

POLYPROPYLENE/GLASS FIBRES INJECTION MOLDED COMPOSITES

ABSTRACT

J. STEIDL, Z. KORINEK, V. ZILVAR

Technical University of Prague - Faculty of Mechanical Engineering Department of Materials Science

Karlovo Nam. 13 - 12135 PRAGUE 2 - Czechoslovakia

The fracture properties of short glass fibre rein­forced polypropylene were studied in relation to micro­structural changes caused by annealing at higher tempe­rature. Two types of injection molded composites were compared. The microstructure was characterized with help of X-ray diffraction and DSC methods. Fracture energy was evaluated on base of notched impact tests, morphology of fracture surfaces by means of scanning electron microscope.

INTRODUCTION

A number of microstructural parameters may influence to a certain extent failure mechanisms of injection mol­ded composites. Therefore, the" structure" approach to the fracture mechanics is of great importance and produ­ces results utilizable in enginnering practice. Micro­structure - fracture behaviour relationships of injecti­on molded composites are systematically studied for exam­ple by K. Friedrich et a1. I TU Hamburg-Harburg I.

In addition to component parts and technological conditions, additional heat exposure influences formati­on of morphology. Some of our previous results indicated that changes in supermolecular structure of matrix may have obvious relevance to the understanding of fracture properties of thermally aged short glass fibre reinfor­ced polypropylene I 1, 2 I. During annealing only micro­structure of matrix is primarily affected, while struc-

288

tural parameters related to fibres remain unchanged. It is the objective of the present paper to point out some observations that are generally known in neat poly­propylene but may result in specific effects providing the material is glass reinforced.

EXPERIMENTAL PROCEDURE

Notched impact test bars were injection molded from two different types of polypropylene filled with 30 % by weight of glass fibres: type A with a good adhesion bet­ween the fibres and matrix I Pr~pathene I and type B with relatively weak coupling I Mosten I.

Supermolecular stbucture of original samples and those annealed at 160 C for 3 hours in vacuum was exami­ned by the use of X-ray diffraction and differential scanning calorimetry.

Fracture energy G was determined on base of impact tests and evaluated witH help of J. G. Williams' method/31 from the relation: GIc = Uc I B W~ , where Uc is the stored elastic strain energy, B is the thickness, W is the width of the specimen and ~ 1S the geometry factor.

Morphology of fracture surfaces was observed in scan­ning electron microscope.

RESULTS

As a consequence of heat treatment, the following changes in the crystalline structure of polypropylene mat­rix were detected:

Transformatton of hexagonal crystal lattice 1/31 to the stable monoclinic~- form I Fig. 1 I. That is to say, certain amount of~- form is being developed during in­jection molding due to shear stresses. ~- form does not appear in quenched samples. Both crystal modifications exhibit apparent texture in surface layers while the inte­rior of specimens was found to be randomly oriented. The ~- form was more pronounced in A-type composite.

The intensity maxima increased much more for A-type composite than for B-type after annealing. The normed diffraction intensity changed from ini~ial value 436 to 1121 I A-type I and from 325 to 648 I B-type I.

An apparent shift in position of X-ray scattering maxima is attributed to disappearing compressive residual stresses in the' surface layers of specimens.

Complicated original DSC curves consisting of at least three melting peaks turn into smooth curve of single mel­ting endotherm I Fig. 2 I.

Shear yielding of matrix surrounding glass fibres is a typical feature of A-type fracture morphology I Fig. 3a I. A good adhesion between matrix and fibres is evident. The annealing process reduces plastic deformation ability

289

of matrix and makes fibre - matrix bonding less strong, as is seen in Fig. 3b. The fracture surfaces of B-type composite have a It patchwork II morphology which signifies a crazing mechanism of failure. This character of fractu­re remained unchanged in heat treated specimens.

The impact fracture energy of A-type composite de­creased markedly as a result of annealing I Fig. 4 I. No appreciable change was observed in case of B-~ype com­posite, Glc is approximately equal to 1, 90 kJm- •

DISCUSSION AND CONCLUSIONS

Heat treatment leds to perfection of crystalline structure and relaxation of residual stresses which is in close relation to mechanical properties of polypropy­lene I 4, 5 I. Resultant phenomena as restriction of seg­mental mobility in amorphous phase, reduction of volume in microregions, shrinkage effects in microvolume, loss of ductility seem to be very important in this respect. It is necessary to emphasize that microstructural chan­ges mentioned above are quantitatively different in neat and in glass fibre reinforced polypropylene I it is a subject of another our study I.

The manifestation of heat treatment varies depending on polypropylene matrix behaviour and fibre - matrix in­teraction. If the matrix is able to deform by shear yielding in the initial state, the microstructural chan­ges caused by annealing substantially influence fracture ener&y. On the other hand, if the matrix inclines to cra­zing mechanism, the changes in semicrystalline structure are probably of no great importance.

Reduction of volume in microregions due to transfor­mation of crystal lattice and higher crystallinity, as well as macro- and microscopic stress relaxation and res­triction of segmental mobility in amorphous phase are thought to be active in debonding process and creation of interspace between fibres and matrix in A-type compo­site annealed at higher temperature. It is to a certain extent consistent with annealing experiments performed on polyphenylenesulphide I 6 I. The authors suggest that debonding forces are mainly caused by stronger mechanical shrinkage stresses of the mor~ crystalline matrix. However, the annealing effect on toughness is rather different for reinforced polypropylene and polyphenylenesulphide. Con­trary to A-type composite, the initial weak bonding in B-type cannot be essentially influenced by microstructural changes in matrix.

The degree of microstructur.al changes also plays an important role in fracture behaviour. It foll~ws fro~ X-ray diffraction experiments that A-type compo~lte.exhl­bits better predisposition to structure reorganlsatlon during annealin~ than B-type.

The anneallng at higher temperatures represents an

290

useful method which enables to reveal some new relation­ships between microstructure and energy consuming mecha­nisms in injection molded engineering composites based on semicrvstalline polymer matrix.

<12,0

Figure 1

dO ""Crt

400

8

2& 20,0>

X-ray diffraction intensity profiles of or1g1-nal I A I, annealed I B I and quenched I C I injection molded specimens of A-type composite. Reflection method, CrK~radiation.

420T [~40 460

Figure 2 DSC curves of original I 1 I and annealed I 2 I A-type composite. Hea~ing range of 16 C min-1•

291

a

b Figure 3 Scanning electron microscope images

of impact fracture surfaces. A-type composite original I a I and annealed I b I.

292

300

-, 200 E

o 50 100 BW, [mm2 ] -----;--

Figure 4 Uc plotted against BW<P for the A-type com­posite. 1 original, Glc = 2, 85 kJm-2

2 - annealed, Glc = 1, 92 kJm-2

REFERENCES

1. Zilvar, V., Korinek, Z. and Steidl, J.: 6th Intern. Conf. on Composite Materials, Imp. College, London 1987. 2. Zilvar, V., Chladkova, Z., Korinek, Z. and Steidl, J.: 2nd Intern. Symp. on Phase Interaction in Composite Materials, Patras 1988. ). Williams, J., G.: Fracture Mechanics of Polymers, Ellis Horwood, Chichester 1984. 4. Greco, R. and Coppola, F.: Plastics and Rubber Pro­cessing and Applications 6 /1986/, 35-41. 5. Coxon, L., D. and White, J., R.: Polymer Eng. Sci. 20 /1980/, 230-236. 6. Karger - Kocsis, J. and Friedrich, K.: J. Mater. Sci. g£ /1987/, 947-961.

COMPATIBILITE CHIMIQUE ENTRE LE MAGNESIUM ET LES FIBRES DE CARBONE

J.C. VIALA, P. FORTIER, G. CLAVEYROLAS, H. VINCENT, J. BOUIX

ABSTRACT

Universite Claude Bernard Lyon 1 Laboratoire de Physico-chimie Minerale

69622 VILLEURBANNE CEDEX - France

PAN-based (DOO) and Pitch-based (P55 and PlOO) carbon fibers were isothermally heat-treated for lOOO-l20h at 450-700 oC in saturated vapour pressure of Ng. Characterization of, the resulting samples by mechanical testing, XRD, SEM, EMA and chemical analysis led to the conclusion that the chemical compatibility of carbon fibers towards Mg depends on their nature pure and highly graphitised Pitch-based fibers exhibit an excellent inertness towards Mg whereas impure and disorded Pan-based fibers may slightly react with this metal.

INTRODUCTION

L'incorporation de fibres de carbone a une matrice a base de magnesium devrait conduire a des materiaux composites possedant une resistance a rupture specifique et surtout une rigidite specifique tres elevees (1). De plus, en utilisant des fibres de carbone a haut module d'elasticite, ces materiaux pourraient avoir un tres faible coefficient de dilatation thermique (1,2). De telles caracteristiques potentielles rendent ces composites particulierement attrayants pour la realisation de pieces destinees a l'aerospatiale, l'aeronautique, voire meme 1 'automobi le et plusieurs programmes de recherChe ont ete developpes dans ce sens (1-3).

Des composites Mg/carbone ont ete elabores en infiltrant sous pression le magnesium liquide dans des preformes carbonees (3-10) ou en compressant a chaud des fibres metallisees (1,2,8). Neanmoins, les essais mecaniques realises sur ces produits ont fait apparaitre des ecarts parfois tres importants entre proprietes mesurees et proprietes theoriquement previsibles par la loi des melanges. Ces ecarts peuvent etre dus a une mauvaise penetration du metal entre les filaments de carbone ou a une repartition trop heterogene de ces filaments dans la

294

matrice. lIs peuvent egalement avoir une orlglne physicochimique telle qu'un defaut de mouillabilite, un manque d'adherence ou au contraire une trop forte interaction chimique a l'interface fibre/matrice. Plusieurs etudes ont ete effectuees dans Ie but d'ameliorer la mouillabilite et l'adherence (2,4-6,8,11-13). Peu de travaux ont en revanche ete consacres au probleme de compatibilite chimique dont nous abordons l'etude dans Ie cadre de ce travail.

I - CONDITIONS EXPERIMENTALES

Af in de prec iser dans que lle mesure des react ions chimiques entre la mat rice et le renfort pouvaient degrader ce dernier lors de l'elaboration a chaud de composites Mg/carbone, nous avons entrepris une etude experimentale a caractere systematique consistant a soumettre des fibres de carbone a des traitements thermiques de longue duree en presence de vapeur saturante de magnesium et a caracteriser les produits obtenus. Cette etude a porte sur trois types de fibres -des fibres T300 a precurseur PAN graphitees a basse temperature -des fibres P55 a precurseur brai assez fortement graphitees ; -des fibres P100 (brai) graphitees a tres haute temperature. Les traitements thermiques ont ete realises en tube de fer scelle sous argon, a quatre temperatures 450°C (lOOOh), 640°C (400h), 653°C (260h) et 700°C (120h). La duree des traitements a ete progressivement reduite ve~s les hautes temperatures pour eviter des reactions parasites avec Ie tube metallique. La caracterisation des produits obtenus a ete effectuee par differentes techniques : d'une part nous avons suivi l'evolution des proprietes mecaniques des fibres traitees en procedant a des essais de traction sur filaments elementaires et en exploitant les resultats en statistique de Weibull (40 a 45 eprouvettes de 20mm de longueur par essai) ; d'autre part nous avons etudie la morphologie, la composition et la microstructure de ces fibres par microscopie electronique a balayage, microsonde electronique, diffraction de rayons X, microsonde Raman et microanalyse chimique.

II - RESULTATS

2.1. Proprietes mecaniques des fibres traitees

Les va leurs de resistance a rupture obtenues sur les trois types de fibres etudies apres traitement thermique en vapeur saturante de Mg sont reportees sur la Fig. 1. Ces valeurs correspondent a une probabilite de rupture egale a 0,5. Dans Ie cas des fibres T300 a precurseur PAN, on observe que la resistance a rupture varie avec la temperature de traitement. Cette resistance, inchangee apres 1000h de chauffage a 450°C en presence de Mg, commence a chuter notablement apres 400h de traitement a 640°C. Cette evolution s'accompagne d'une diminution du parametre de Weibull traduisant une plus grande dispersion des resultats. Ces tendances s'accentuent lorsque les fibres ont ete traitees 240h a 653°C: a ce stade, elles ont en effet perdu environ la moitie de leur resistance a rupture nominale alors que des fibres chauffees dans les meme conditions en l'absence de Mg n'ont subi pratiquement aucun dommage. Enfin, apres 120h de traitement

295

thermique a 7000 C dans Mg, on retrouve une resistance a rupture inferieure a la valeur nominale, mais nettement plus elevee qu'a 653°C. Cette evolution de proprietes mecaniques presentee par les fibres T300 n'est pas observee dans Ie cas des fibres P55 ou PlOO, pour lesqelles la resistance a rupture reste tres proche de la valeur nominale, quelle que soit la temperature de traitement dans Mg.

2.2. Microstructure, morphologie et composition

Des informations sur la microstructure des fibres brutes et traitees ont ete obtenues en determinant systematiquement, par diffraction de rayons X, Ie parametre de maille moyen c des cristallites de graphite constituant les fibres et l'extension moyenne Lc de ces cristallites dans la direction de 1 'axe c (Fig. 2). 11 apparait que les fibres T300 subissent une evolution microstructurale par traitement thermique en vapeur de Mg : le parametre c diminue regulierement quand la temperature croit de 450 a 7000 c alors que la valeur de Lc augmente. Cette evolution, qui ne se produit pas en l' absence de Mg, correspond a un rearrangement des feui llet s graphitiques au sein des fibres T300. Elle n'est observee ni avec les fibres P55, ni avec les fibres PlOO soumises au memes traitements.

L'examen systematique par MEB des fibres brutes et traitees n'a pas permis de mettre en evidence d'evolution importante de la morphologie de surface due au traitement en presence de Mg.

Les teneurs massiques en C, N et 0 des fibres brutes ont ete determinees par microanalyse chimique. Les resultats suivants ont ete obtenus : - fibres T300 : C = 91,8%, N = 8,1%, 0 = 0,1% j

- fibres P55 : C = 99,2% N 0,02%, 0 = 0,2% j

- fibres PlOO : C = 99,6%, N = 0,02%, 0 = 0,1%. Apres traitement thermique dans Mg, la composition des fibres P55 et PlOO est restee inchangee, alors que dans Ie cas des fibres T300, une diminution de la teneur en azote a ete observee apres traitement a 640, 653 et 7000 C (environ 6,4% au lieu de 8,1% initiaiement). Des analyses effectuees par microsonde electronique sur la section de filaments ont .confirme la baisse sensible de la teneur en azote des fibres T300 apres traitement. Elles ont de plus montre que ces fibres pouvaient contenir jusqu'a 20% massiques de Mg, reparti dans Ie volume de chaque filament. Des analyses complementaires effectuees par microsonde Raman ont en outre montre que ce magnesium se trouvait essentiellement a l'etat metallique. Du magnesium a egalement ete observe a l'interieur des filaments de fibres P55, mais en proportion dix a vingt fois moin.dre, alors que cet element n'a pas ete detecte dans les fibres P100. Enfin, i1 est a noter que ni Mg, ni aucune des phases telies que MgO, Mg3N2, MgC 2 ou Mg 2C, n'a pu etre mise en evidence par diffraction de rayons X dans les fibres traitees.

III - DISCUSSION

Les resultats precedents font apparaitre que les fibres de carbone a precurseur brai de type P55 ou PlOO ne subissent pratiquement aucune alteration de proprietes mecaniques, de microstructure et de composition lors de traitements thermiques

296

prolonges en presence de Mg. On peut donc conclure que ces fibres presentent une excellente inertie chimique vis-a-vis de ce metal, dans le domaine de temperature 450-700 0 C.

En revanche, traitees dans les meme conditions, les fibres T300 a precurseur PAN perdent une partie de leur resistance a rupture, tendent a recristalliser et evoluent en composition. Ces fibres sont donc chimiquement moins inertes vis-a-vis de Mg que les precedentes. On peut penser que cette difference de comportement resulte du fait que les fibres PAN sont constituees d'un graphite moins dense, beaucoup moins pur et nettement plus desordonne que les fibres Pitch; Neanmoins, la reactivite des fibres PAN vis-a-vis de Mg est toute relative puisque, bien que des traitement particulierement longs aient ete effectues, les proprietes mecaniques, la microstructure et la composition des fibres ont peu varie et aucun produit de reaction n'a pu etre caracterise par les techniques mises en oeuvre. On en est donc reduit a faire des hypotheses sur les interactions susceptibles de se produire entre Mg et les fibres T300.

Une premiere possibilite est la formation de composes d'insertion tels que ceux observes avec les metaux alcalins (14). Une autre eventualite est la formation de carbures de magnesium, MgC 2 ou Mg 2C1 , par combinaison directe des elements : bien que la synthese directe oe ces carbures n'ait jamais ete clairement mise en evidence, celle-ci reste en effet thermodynamiquement possible aux temperatures i nferieure s a 485°C pour MgC 2 (15) et a 670°C pour Mg 2C3 (16). Enf in, on peut prevoir des interactions entre entre Mg et l'azote ou l' oxygene contenus en impuretes, avec format ion d' oxyde MgO ou de nitrure Mg3N).

L'hypotfiese de la formation d'un compose d'insertion ne semble pas devoir etre retenue. En effet, l'insertion d'elements electropositifs dans le graphite se traduit toujours par une forte dilatation du parametre de maille c, or c'est a une contraction que l'on assiste. En outre, si on trouve des quantites relativement importantes de Mg ai' interieur des filaments traites, celui-ci est essentiellement a l'etat metallique non combine. Pour cette derniere raison, la formation de quantites appreciables de carbures parait egalement a exclure. Reste l'hypothese de reactions entre Mg et l'azote ou l'oxygene contenus en impuretes dans les fibres T300 : en ce qui concerne l' oxygene, la teneur toujours tres basse de cet element n'a pas permis d'observations j dans le cas de l'azote, les variations de teneur mises en evidence tendent a prouver qu'une reaction chimique se produit effectivement.

Quant a la presence de Mg a l'etat metallique dans les fibres traitees, elle ne peut s'expliquer que par un processus de condensation capillaire intervenant dans les pores des fibres T300 au cours du traitement en vapeur saturante de Mg. Ce phenomene ne concernant que les pores de tres petit diametre (moins de 500AO), les cristaux de Mg ainsi formes ne peuvent eux-meme avoir qu 'une tres petite taille, ce qui expliquerait que, bien qu'ils soient relativement abondants, ils n'aient pas pu etre caracterises par diffraction de rayons X. Ce phenomene s'observerait egalement avec les fibres pj5, mais a un degre bien moindre.

La condensation capillaire etant un phenomene physique limite a la surface des fibres, il est tres peu probable qu'elle affecte leur

297

microstructure ou leur propri~t~s m~caniques. En revanche, Ie d~part d'une fraction de l'azote initial par r~action avec Mg est tres vraisemblablement a l'origine de la recristallisation des fibres T300. En effet, la pr~sence d'azote dans ces fibres peu graphit~es est un facteur de d~sordre, cet ~l~ment se trouvant engag~ dans des h~t~rocycles carbone-azote. La destruction partielle de ces h~t~rocycles ne peut donc que favoriser un r~arrangement au sein des feuillets graphitiques, en activant la diffusion d'atomes de carbone a basse temp~rature. La diminution de r~sistance a rupture observ~e apres traitement des fibres T300 serait la cons~quence de cette recristallisation. N~anmoins, l'~volution de microstructure des fibres T300 se fai sant de facson monotone avec la temp~rature, la recristallisation ne suffit pas a justifier Ie fait que la r~sistance a rupture des fibres T300 passe par un mlnlmum vers 653°C et l'hypothese d'une r~action directe entre Mg et Ie graphite tres d~sordonn~, de ces fibres, avec production de traces de carbure Mg 2C3 , n'est pas a ~carter.

CONCLUSION

Les r~sultats de cette ~tude mont rent que la compatibilit~ chimique entre le magn~sium et les fibres de carbone d~pend de la nature de ces dernieres. Des fibres relativement pures et fortement graphit~es, telles que les fibres P55 ou P100 a pr~curseur brai, pr~sentent une excellente inertie chimique vis-a-vis de ce m~tal, dans Ie domaine de temp~rature 450-700oC. En revanche, des fibres impures constitu~es d'un graphite tres d~sordonn~ telles que les fibres T300 a pr~curseur PAN peuvent donner 1 ieu a des ph~nomenes d' interact ion chimique avec Ie magn~sium : ces ph~nomenes restent n~anmoins tres limit~s et ne peuvent suffir a expliquer les deviations importantes a la loi des m~langes report~es pour certains composites Mg/carbone.

REFERENCES.

1 • Ph.ROY,A. MAMODE, Proc. 3rd Europ. Symp. on Spacecraft Materials in Space Environement, Noordwijk (N.L.), Oct. 1985, 185-90. 2 • O. REMONDIERE, R. PAILLER, A.MAMODE, Ph ROY, Proc. ECCM1, Bordeaux (Fr.), Sept. 1985, 732-37. 3 • D.M. GODDARD, W.R. WHITTMAN, R.L. PUMPHREY, Proc. Int. Congo SAE, Detroit (USA), Feb. 1986, 3598-603. 4 A.P. LEVITT, E. Di CESARE, S.M. WOLF, Met. Trans, 3 (1972) 2455-9. 5 M.H. RICHMAN, A.P. LEVI6T, E. Di CESARE, Rep. 754570, NTIS (1972). 6 H.A. KATZMAN, US Pat. n 4 376 804, 15 Mar. 1983. 7 S.I. DEMENTEV, A.A. ZABOLOTSKII, I.V. ROMANOVICH, S.A. PROKOFEV, S.E. SALIBEKOV, Poroshk. Metall. (Kiev),3 (1977) 50-4. 8 • S.P. RAWAL, L.F. ALLARD, M.S. MISRA, Proc. ICCM6, London (G.B.), Jul. 1987, 2169-82. 9 • A.P. DIWANJI, I.W. HALL, ibid., 2265-74. 10. G.D. LAWRENCE, Trans. Am. Foundrymen's Soc., 80 (1972) 28,3-86. 11. Y. NAERHEIM, M.W. KENDIG, 18th SAMPE, Seatle (USA), Oct. 1986, 12. H. KATZMAN, Report n SD-TR-88-63 ,NTIS (1986). 13. H.A. KATZMAN, J. Mater. Sci., 22 (1987) 144-48. 14. W. RUDORFF, Chimia, 19 (1965) 489-99.

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15. B. HAJEK, P. KAREN, V. BROZEK, Collect. CZE'ch. Chern. Commun. 48 (1983) 1963-68. 16. B. HAJEK, P. KAREN, V. BROZEK, ibid., 48 (198:1) 196:1-b8.

.. (0 • . .... , ... U"

• 'I"

(a) PAN T300

--r-

.... UI ... R c-

'" T( e)

.. 00' .... ....

(b) PITCH P55 (c) PITCH PlOO

,,,. -

r( C) fC"C)

Fig. 1. Resistance 11 rupture (Pr = 0,5) de fibres de carbone apres traitement thermique en vapeur saturante de magnesium.

(a) PAN T300 (b) PITCH P55 (c) PITCH PlOO

c

" (l) (11 ( ll

"

.. """

~.: ~ ~.,

~, 'I' f"C I f tC)

< ., " ClI (.I. ) (.I.)

...

... T("C)

Fig. 2. Parametre de maille c et taille des cristallites de graphite Lc de fibres de carbone traitees en vapeur saturante de magnesium.

CONCEPTION ET CALCUL DESIGN AND ANALYSIS

Chairmen: Pr F. L. MATTHEWS Imperial College of Science and Technology PrS. W. TSAI US Air Force Dr J. A. N. SCOTT Shell Laboratorium

ETUDE DE LA FISSURATION D'UN MATERIAU COMPOSITE VERRE-EPOXYDE UNIDIRECTIONNEL SOLLICITE EN TRACTION

G. PLUVINAGE, C. SCHMITr, M. ABISROR"

Universite de Metz - Faculte des Sciences /Ie du Saulcy - Laboratoire de Fiabilite Mecanique - 57045 METZ - France

'Institut de Soudure

ABSTRACT

Zone Industrielle des Jonqu;eres - 57640 VIGY - France "Aerospatiale - Laboratoire Central

12 rue Pasteur - BP 76 - 92152 SURESNES - France

Fracture experiment were carried out on unidirectional glass epoxy composite material with the presence of a crack parallel to the fiber direction. Fatigue tests were performed using tensile split specimens, the crack being parallel to the fiber direction. Split specimens loaded In fatigue tension parallel to the fibers were used. A damaged zone consisting in fiber bridging the major crack was observed and micro fracto graphic observations revealed that the crack extension lad to mode one fracture mechanism. A three dimensional finite element method was employed to calculate the values of the strain energy release rates.

Crack growth rates versus ~K 1 and ~K2 were studied and compared to those of pure mode I and mode II obtained with DCB specimens. It was shown, that this "analysis" corrolate microfractographlc observations.

INTRODUCTION

Dans Ie cas de maU:riaux anisotropes, la Mecanique Lineaire Elastique de la Rupture ne peut pas etre appliquee de fat;on generale comme aux materiaux isotropes. En effet, il est necessalre d'apporter certaines restrictions que Paris et Sih 11/ ont precise :

- la direction initiale de la fissure dolt coincider avec un axe de symetrie elastique.

- la fissure dolt se propager dans son plan et sa direction initiale.

Dans ce cas particulier, il n'existe pas de couplage mecanique des modes de base et n est possible de deflnir les facteurs d'intenslte de contraintes de fat;on similaire a ceux des corps isotropes, gouvemant la meme Singularlte en 1/ ~r en fond de fissure 12/. II existe alors une relation blunlvoque entre facteurs d'intensite de contraintes et taux de restitution de l'energle. Dans Ie cas de l'orthotrople, on a en contralntes planes :

302

011 les coefficients Sij sont les coefficients de la matrice de complaisance elastique.

Dans Ie cas de materiaux composites verre-epoxyde unidirectionnels comportant une fissure parallele aux fibres, les conditions decrites ci-dessus sont touJours respectees. En effet, la fissure etant parallele aux fibres, elle coIncide avec un axe de symetrie elastique.

De plus, l'experience montre que la fissure se propage dans son plan et dans sa direction initiale qui est celle des fibres, cela aussi bien en mode 1 [3, 4, 5, 6 ] qu'en mode 2 [ 6, 7, 8 ].

I - REALISATION EXPERIMENTALE

1.1. Materiau et eprouvettes

Un materiau composite verre-epoxyde unidirectionnel monocouche a ete utilise, la fraction volumique . de fibres etant de 54 %. Les constantes elastiques ont ete determinees par des essais classtques de traction et cisaillement :

52000 MPa 15300 MPa

(direction des fibres)

V 12 = 0,3 (V12 = £2/£1)

G12 = 3 800 MPa

Les eprouvettes ont ete decoupees dans des plaques unidirectionnelles obtenues en disposant des preimpregnes en quincortce en autoclave. Un clinquant d'acier a ete introduit pour initier un defaut Ie long des fibres. Des eprouvettes "echardes" ont ete obtenues en donnant un coup de scie perpendiculaire au defaut. Des talons en aluminium colles sur chaque face permettaient Ie maintien de l'eprouvette dans les mors de la machine d'essai (figure 1).

1.2. Methode experimentale

Des essais de fatigue ont ete realises en traction a charge tmposee sur une machine de traction. Les suivis de fissure ont ete realise a l'aide de lunettes grosstssantes (30 fois). Apres rupture, les facies de rupture etaient examines a l'aide d'un microscope electronique a balayage.

Pour comparer les resultats experimentaux et les calculs par elements finis, les deplacements relatifs des levres de la fissure ont ete mesures a l'aide d'un capteur a lames. Celui-ci etait maintenu par deux pieces en aluminium collees sur l'eprouvette de part et d'autre de la fissure(figure2).

1.3. Modele analytique

Le modele de wagner 191 applique a la geometrie de l'eprouvette echarde a ete utilise pour evaluer Ie taux de restitution de l'energle global et compare aux previsions des calculs par elements finiS.

Selon ce modele. Ie taux de restitution de l'energie est donne par la formule:

G = FC2 . b/2El e2 I (I-b)

ou Fc : force appliquee, I : largeur de l'eprouvette, b : largeur de l'echarde

303

Cette expression est identique a celle qu'on aurait pour un matertau isotrope du fait de l'ortentation de la charge par rapport aux axes de symetrie.

Pour Ies faibles Iargeurs d'echarde (1- b _ I), G est proportionnel a Ia largeur d'echarde b.

350 I - •

t I , , -1 b : largeur .~e

I echarde - --- -- --- --_._------ ---- -- _. --f--

'II ~ sens des f ibres -,

" 1 ' .. .! ,,. • !

65 5 5 65

traction 4j tract ion

+- ===~::=======f-====::::::::::===' .... " talons en aluminium

Fig. 1 - Eprouvette echarde

Cette expression est identique a celle qu'on aurait pour un matertau isotrope du fait de l'ortentation de la charge par rapport aux axes de symetrie.

Pour Ies faibles Iargeurs d'echarde (1- b _ I), G est proportionnel a la largeur d'echarde b.

plots de mesure

~'d".'P'.~ lames du capteur j \ cor

Fig. 2 - Methode expertmentale de mesure des deplacements

II - SIMULATION NUMERIQUE PAR ELEMENTS FINIS

2.1. Methode numertque

Les calculs par elements finis ont ete realises en trois dimensions a l'aide du code de calcul Samcef, developpe par l'Universite de Liege, en utilisant des elements multicouches a interpolation Uneaire modeUsant un matertau composite stratlfie. Pour chaque eprouvette, Ie maillage contenait 880 elements, dont quatre rangees dans l'epaisseur.

Une methode inspiree de Ia methode de Rybicki 1101 et developpee l'Aerospatiale a permis de determiner la partition des modes de rupture, c'est-a-dire les valeurs respectives de G, Gl, G2 et G3.

304

2.2. Resultats numertques

- Les valeurs du taux de restitution de l'energie sont independantes de la longueur de fissure. De plus, les calculs donnent un resultat conforme aux previsions du modele de Wagner (G theortque),(figure 2).

- La partition des modes est lndependante de la longueur de fissure, mais varte legerement avec la quantite b, largeur de l'echarde (tableau 1).

- Le mode 3 est en fait negligeable. Son existence est due Ii l'effet de striction de l'eprouvette et peut etre expliquee en examinant les valeurs de G3 Ie long du front de fissure (tableau 2).

b (mm) G1 (0/0) G2 (0/0) G3

4 38,4 a 39,1 59,3 a 60,0 1,5

5 37,7 Ii 38,1 60,3 Ii 60,7 1,5

6 36,7 Ii 37,0 61,3 a 61,7 1,6 6,5 35,9 a 36,7 61,7 Ii 62,5 1,6

Tableau 1 : Variation de la partition des modes en fonction de la largeur de l'echarde b

Front de fissure bord milieu bord

G1 (J/m2) 467 550 567,5 550 467

G2 (J/m2) 882 812 831 812 882

G3 (J/m2) 75 6,3 0 3 75

Tableau 2 : Variation de G3 Ie long du front de fissure (a = 40 mm, b = 4 mm, Fc = 4000 daN)

III - RESULTATS EXPERlMENTAUX ET DISCUSSION

3.1. Mecanisme de propagation

(0/0)

Lors des essais, n. est possible d'observer des paquets de fibres non rompues en travers de la fissure et jOignant les deux levres, en den site moindre qu'en mode 1 pur 14, 9/.

Les observations fractographiques montrent que globalement les facies s'apparentent Ii des facies de rupture en mode I : la rupture se produit par cl1vage dans la resine.

Cependant, on peut trouver localement des languettes microscopiques dans la serte inclinees par rapport a la direction des fibres caractertstiques du mode II ou du mode mixte 16/.

305

La part predominante du mode I de rupture dans ces mecanismes de fissuration n'est pas surprenante si 1'0n considere la valeur du rapport G2c/Glc pour Ie materiau utilise:

G2c/GIC _ 330 J/m2/3 300 J/m2 _ 10

3.2. Mesure des deplacements

Sur une courbe typique charge deplacements relatifs des levres de la fissure, n est a remarquer que les deplacements calcules sont legerement superieurs a ceux mesures, probablement a cause de la presence de la zone endommagee, tendant a accroitre la rigidite de l'eprouvette.

La modelisation par elements finis ne permettant pas de prendre en compte la zone endommagee, les depouillements ulterieurs constitueront une prevision conservative de la tenue en fissuration du materiau.

3.3. Vitesse de fissuration

Les vitesses de fissuration ont ete exprimees respectivement en fonction

de L\kl et L\k2, les facteurs d'intensite de contraintes etant obtenues a partir des valeurs du taux de restitution de l'energie calculees directement par elements finiS. Les resultats regroupes sur la figure 3 montrent que la fissuration est gouvernee par la loi de Paris. Dans cette analyse, les effets de la fissuration stable (fissuration sous tension), d'ailleurs negligeable, n'ont pas ete pris en compte 14/. La comparaison des essais de traction d'echardes avec des essais de mode 1 et mode 2 purs realises a l'aide d'eprouvettes DCB met en evidence les pOints suivants :

- En mode 1, on obtient une droite parallele aux essais a deplacement impose et on reJoint pour les faibles vitesses de fissuration Ie nuage de pOints correspondant aux zones moyennement endommagees a force imposee. II est important de rappeler qu'en mode 1 pur a force imposee, au fur et a mesure que la fissure avancait (done Kl augmentait), la zone endommagee s'amplifiait et on obtenait par consequent une vitesse de propagation decroissante avec

L\Kl 113/.

- En mode 2, la comparaison met en evidence deux droites paralleles. Vu ces analyses et les exam ens fractographiques, 11 est legitime de

supposer non pas une propagation en mode mlxte 1 + 2, mais une propagation en mode 1 pur uniquement, Ie mode 2 prenant part dans la destruction partlelle de la zone endommagee, puisqu'en mode 1 pur, la zone endommagee est tres abondante, alors qu'en mode 2 pur, elle est inexistante.

Ainsi, bien qu'il n'existe pas de couplage mecanique des deux modes de rupture, les conditions definies precedemment etant respectees (direction de la fissure confondue avec un axe principal de symetrie se propageant dans son plan initial), 11 existe un couplage physique des deux modes.

Par consequent, Ie mode 2 represente lei un facteur accelerant. Ce phenomene n'est pas rencontre chez les materiaux metall1ques soll1cites en mode mixte, car la fissure bifurque pour prendre une direction telle qu'elle puisse se propager en mode 1.

306

i (J > (J ...... E g z 0 ...... 4 0

• _2

10

-3 10

_4 10

•

R: 0,3

2 J 4 5 6 .1K(MPaVm)

Fig. 3 - Vitesse de fissuration en fonction de ~K 1 et ~K2 pour les eprouvettes echardes

CONCLUSION

Cette etude a mis en ev1dence les mecanismes de rupture d'un matertau composite verre-epoxyde unidirectlonnel soll1cite en traction a l'aide d'eprouvettes echardes. Bien que les soll1citaUons engendrent des composantes de mode mtxte, 11 a He montre que seul Ie mode 1 prend part dans la propagation de la fissure. Quant au mode 2. bien que n'intervenant pas proprement dlt dans la propagat1on. 11 reste cependant present en contribuant a une destruction partielle de la zone endommagee constltuee de fibres en travers de la fissure.

Les memes phenomenes ont Ne observes sur Ie meme materiau sollicite en traction-torsion. Cette etude a egalement montre la puissance des elements finiS. indispensables pour traiter ce genre de probleme. Cependant, Ie fait de ne pas tenir compte de la presence de la zone endommagee dans la model1sation fait que les resultats presentes sont une prev1sion conservatlve de la tenue en fissuration du materiau.

REFERENCES

1 - PARIS P.C and SIH G.C .. ASTM STP 381 ; p. 30-81. (1965)

2 - SIH G.C .. PARIS P.C .. IRWIN G.R.. IJF Vol. 1. W 3. p. 189-203. (1965)

3 - WU E.M .. Journal of applied Mechanics, p. 967-974. (1967)

4 - SLEPETZ J.M and CARLSON L., ASTM STP 593, p. 143-162. (1975)

5 - PARHIZGAR S .• iowa State University. PHD (1979)

6 - NESA D .• These presentee a l'ecole des Mines de Paris, (1987)

307

7 - WU E.M .• Fracture and Fatigue. Composite Materials. Vol. 5. (1974)

8 - G.S. GlARE. Eng. Fract. Mech .. Vol. 20, n° 1, p, 11-24 (1984)

9 - WAGNER D. J. of Material SCience, Letters nO 5, p. 229-230. (1986)

10 - RYBICKI E.F. and F. KANNINEN. Eng. Fract. Mech .• Vol. 9. p.931-938. (1977)

11 - SCHMITT C .• These Nancy, (1988)

DESIGN AND ANALYSIS OF ORTHOTROPIC COMPOSITE MATERIALS THROUGH A MIXED MODE

COHESIVE CRACK SIMULA nON

ABSTRACT

S. VALENTE. A. CARPINTERI

Politecnico di Torino - Department of Structural Engineering 10129 TORINO -Italy

A Mixed Mode crack analysis is applied to simulate the failure mechanism of the Iosipescu shear test for a fibre-reinforced composite. The stability of the curvilinear crack trajectory appears to be sensitive to the ratio between the principal strengths. For a strongly orthotropic material, the zig-zag trajectory results to be alternately parallel to the fibres or directed towards the loading poi nt.

- INTRODUCTION

Objective of this work is that of developing the capability to analyze accurately the mixed mode propagation of a crack /1/ in an arbitrary structure with elastic orthotropic material stiffness properties and anisotropic material strength characteristics. Whereas for Mode I, only finite element node untieing is applied to simulate crack propagation, for Mixed Mode interelement crack propagation a continuous modification of the mesh is required /2/. The numerical response of a Iosipescu shear specimen /3/ is presented using a computer simulation founded on a Mixed Mode cohesive crack model /4,5/. At each computing step the fictitions crack tip propagates by a pre-defined length A a in a direction orthogonal to the maximum normalized circumferential stress.

The stability of the curvilinear crack trajectory is analyzed by

310

varying the ratio between the principal strengths or toughnesses. Whereas for an isotropic material the trajectory is smooth and connects regularly crack tip and opposite force, for a ~trongly

orthotropic material the trajectory firstly appears as orthogonal to the direction of minimum strength (i.e., parallel to the fibres) and then, suddenly, the crack changes its path to run towards the opposite force. In this latter case, therefore, a remarkable instability of the trajectory is evident.

With the same numerical procedure it is possible to study Mixed Mode crack propagations also at the interfaces between different orthotropic materials or laminae.

2 - MICRO-MECHANICS OF FIBRE PULL-OUT

The assumptions used to study Mixed Mode crack propagation in an orthotropic material, will be proposed and explained in the present section.

The orthotropic strength is a vectorial property depending on the fibre direction. It is proportional to the fibre strength and the fibre percentage, the matrix being usually very weak in comparison with the reinforcement. If the fibres are aligned, the percentage in volume is egual to the percentage on the orthogonal cross-section (Fig.1-a):

On the other hand, the percentage in volume is connected with the percentage on an inclined cross-section by the following relationship (Fig.1-b):

which is reduced to: (1-c)

for a cross-section parallel to the fibres (Fig.1-c). If we cons i der a fi bre of radi us R, inherent strength ° f and

adhesion shear strength "it is possible to determine the maximum 1 ength of s 1 i ppage 1 : s

s

2 11 R 1 , 2 = 11 R °

Eq. (2) is valid only if the fibre-matrix interface.

1 s

s s f

a uniform shear From eq. (2) it

RO f /2's

(2)

stress, is assumed at results:

311

When 1 > 1 ,the tensile failure of the fibre precedes its slippage (Fig.2'. On the other hand, when 1 < 1 , the slippage

s strength is:

a = 2 '[ 1 1 R < of' S s

(4)

If the fibre bridging a crack is supposed to be long 21, the average slippage strength is:

1 1 - = { as d1 1 f d1, a

S

o 1 o 1 1

a = ( f So s d1 + I of d1) 1 f d1, s 0 1 0

s Recall ing eq. (4), eqs. ( 5) become:

as = 1/2 of 1 1\ ' as = of (1 - 1/2 . 1 /1 ) ,

-

for 1 < 1 , - S

for 1 > 1 • s

for 1 ~ 1 , s

for 1 > 1 s

(5-a)

(5-b)

(6-a)

(6-b)

The relationship a versus 1 is plotted in Fig.3. It is linear s between zero and 1, and then hyperbolic with an horizontal

s asymptote:

1 ima 1 + '" S

= a f

(7)

When the fibres are aligned and the fibre length is large, 1> 1 , the application of eqs (1-b) and (6-b) provides the ultimate tensite strength of the fibre-reinforced material:

a being the tensile strength of the matrix, while the critical v~lue w of the crack opening displacement provides the fracture energy~\c through the rellati1~)nship (Fig.4):

~IC = 1/2 au Wc • (9)

In the next sections, a composite material with long and aligned fibres will be analyzed, with the same orthotropic variation of tensile strength and fracture energy as functions of the angle ~ -see eqs (8) and (9) - w being assumed as an isotropic property.

c

312

Similar variations are proposed in the literature, as, for example, in 16/.

3 - MIXED MODE COHESIVE CRACK PROPAGATION

The propagation criterion is plotted in Fig.5. The external locus is the polar representation of the orthotropic strength, eq.8, while the internal one is the polar representation of the normal stress.­The direction at the tangential point provides the predicted direction of crack branching.

According to the "fictitious crack length control scheme" 14/, the loading process is controlled by the crack length, which is alwa~s a monotonic increasing function of time.

4 - DISCUSSION

The numerical response of the Iosipescu shear specimen (Fig.6) is analyzed according to the cohesive crack model 14,5/.

The geometrical features of the specimen are the following: 1 3.5 b, a = 0.25 b, c = 0.864 b, and the material is assumed to present o~thotropic properties: E1 = 3.45 E2, G12 = 0.244 E2, v =

0.1, w = 0.02 b. The shape of the craCK trajectories changes substa~tially by varying the orthotropic strength ratio a /a 2 =

qICl/~IC2 and keeping the geometrical shape of the ustrtlcture unchange~. It can be observed (Fig.7) that the crack trajectory is smooth when the strength ratio is equal to or less than 5. For larger values, the crack shows some sudden branching in the fibre direction and, alternately, towards the loading point.

Eventually, when the strength ratio is equal to or greater than 50 the crack propagation is collinear to the fibre direction.

These considerations on the crack trajectories are in agreement with the experimental results obtained by Kumosa and Hull 171.

ACKNOWLEDGEMENTS

The financial support of the Department of Public Education (M.P.I.) is gratefully acknowledged.

tIl ! ! ! I" tIt

I • • j II" I IIlIII (a) (b) (c )

Fig. 1 - Fibre orientation

'" f-o

~ rr w 0: S f-o <Il

'" d f ~ "" E; -' til

'" ~ ~ 1 <

0 1 s

FIBRE HALF - LENGTH Fig. 3 - Average slippage strength

vs. fibre half-length

Cip seress funcCion

y

313

4 _ cl /(:'.'?":i::% s

d -:/,../~. - elf f~ ///:

/ , ~. /,,. .... 1

_, _ 5_

(a) (b)

Fig. 2 - Slippage and tensil e failure of a fibre

d

d <Il U <Il W 0: f-o <Il

W o Wc

CRACK OPENING DlSPLACEME~7

Fig. 4 - Stress vs. crack opening displacement

Fig. 5 - Fracture propagation cri­terion for an orthotropic materi al

Fig. 6 - Iosipescu shear test

314

D

E ~ F .-..3.(L

G ~

H 50

Fig. 7 - The effect of the orthotropic strength ratio 0Ul/oU2 on the crack trajectories

REFERENCES

1 - G.C.Sih, P.C.Paris and G.R.Irwin, "On cracks in rectilinearly anisotropic bodies", Int.J.Fracture, 1 (1965) 189-203.

2 - T.J.Boone, P.A.Wawrzynek and A.R.Ingraffea, "Finite element mode 11 i ng of fracture propagati on in orthotropi c materi a 1 s", Engineering Fracture Mechanics, 26-2 (1987) 185-201.

3 - N.Iosipescu, "New accurate procedures for single shear testing of metals", Journal of Materials, 2 (1967) 537-566.

4 - A.Carpinteri and S.Va1ente, "Size-scale transition from ductile to brittle failure: a dimensional analysis approach", CNRS-NSF Workshop on Strain Localization and Size Effect due to Cracking and Damage, (Sept.6-9, 1988) Cachan (France).

5 - A.Carpinteri, S.Va1ente and P.Bocca, "Mixed mode cohesive crack propagation", Seventh International Conference on Fracture, (March 20-24, 1989) Houston (Texas, USA).

6 - A.M. Brandt, "Influence of the fibre orientation on the energy absorption at fracture of SFRC specimens", Euromech Colloquium 204, Structure and Crack Propagation in Brittle Matrix Composite Materials, (A.M.Brandt and I.H.Marsha11, eds.) (1986) 403-420, Elsevier, Jablonna (Poland).

7 - M.Kumosa and D.Hu11, "Mixed-mode fracture of composites using Iosipescu shear test", Forth Int. Conference on numerical methods in fracture mechanics, (1987) 657-667, Pineridge Press, San Antonio (Texas, USA).

CONTACT BEHAVIOURS OF LAMINATED COMPOSITE THIN SHELLS AND A RIGID BALL

L.J. LEE, C.K. PAl', L.C. SHIAU

Institute of Aeronautics and Astronautics National Cheng Kung University

70101 TAINAN TAIWAN - Republic of China 'Chung-Shan Research Institute

PO Box 1-26-2 Lung-Tan, Tau Yvan TAIWAN - Republic of China

ABSTRACT

Static indentation laws for laminated comp(l~ite thin shells and a rigid ball are proposed in this research. The contact behaviors of the rigid ball and target shells are divided into three stages i.e., the loading, unloading and reloading stages. While a 2.5 power is suggested to use in unloading stage, it is found that Hertzian type 1.5 power law is adequate to represent the loading and reloading paths. The effect of the curvature is considered by imposing an unknown function in force-indentation equation for describing the contact behavior at loading stage. This unknown function which characterizes the curvature effect is obtained by modifying the Hertzian law. Once this unknown function is determined, the unloading and reloading equations could be derived from it.

INTRODUCTION

Recently, the fiber reinforced laminated composites have been used success­fully to replace the conventional aluminum alloy to build the skin of the wing and fuselage of an aircraft. This is due to the fact that the former has higher specific modulus and strengths than the latter. But owing to the lack of transverse rein­forcement, the laminated composites are susceptible to damage when subjected to impact loading. For the past decade, a lot of reearches have been undertaking to improve the damage resistence of laminated composites subjected to impact force [1,2]. In order to quantify the impact damage accurately, it is necessary to calculate the impact force and responses precisely. In most existing analytical studies [3,4], the impact loading is assumed to be prescribed. This assumption could simplify the analysis. However, the dynamical responses of the composites might not agree with the experimental results.

Since the laminated composites are anisotropic and nonhomogeneous, the contact behavior between an impactor and laminated composites is very compli­cated. To derive an analytical form of contact behavior of laminted composites

316

is not only very complex but also unrealistic. To circumvent this difficulty, Yang and Sun [5] used an experimentally determined static indentation law to describe the contact behavior. Subsequently, many researches [6 -8] on the impact re­sponses of laminated composites were accomplished based on the aforementioned indentation law. In reference 5, the proposed laws for laminated plate consist of three stages, i.e., the loading, unloading and reloading stages. They describe the contact behavior of a rigid spherical impactor and laminated composite plates. Lee and Huang [8] employed these laws in the analysis of impact responses of a laminated cylindrical shell.

In order to calibrate the contact force as accurately as possible for using in the impact damage analysis of composite shells, contact laws for shells should be established first. Because of the nature of geometry, contact behavior for compos­ite shells are different from that for plates. The effect of curvature of a composite shell should be included in the contact process. An experimental procedure sim­ilar to that of reference [5] is adopted herein to establilsh the contact laws for rigid balls and composite thin shells which are manufactured by graphite/epoxy prepreg .. A mathematical modeling of the loading path is also proposed.

I - SPECIMEN DESCRIPTION AND TEST SETUP

The composite thin shells were made of graphite/epoxy which is manufac­tured by Industrial Research Institute of Taiwan. The laminar properties are given as follows:E1 = 137.9GPa, E2 = 10.34 GPa, G12 = 6.55 GPaandl/12 = 0.21. The prepreg was laid up on a aluminum cylindrical mold. Two kinds of stacking sequence, i.e., [0°/45% °/ - 45° /ooh. and [90° /45° /90° / - 45° /90°12. are used in this research.

The 90° fiber is oriented along the longitudinal axis of the cylinder. The layup set was then cured in an autoclave to form a cylindrical thin shell with nominal thickness of 2.77 mm. The cured shells were cut out by a diamond blade machine to obtain 3 cm wide specimen. There are four different radii of curvature (Rd, i.e., 10 cm, 15 cm, 23 cm and 39 cm for the curved specimens.

The indentation test setup was designed similar to that in Ref. [5]. The indentation was measured by a COMPAC 555 dial gage whose minimum reading is 0.002 mm. The dial gage was mounted on a "C" bracket fixed to the load cell. Only the relative deformation of the indentor and the specimen was recorded. In all tests, the specimens were clamped at both ends and its span is fixed to be 7.5 cm. Three steel balls of diameters 6.35 mm, 12.7 mm and 25.4 mm were used as indentors. This setup was mounted on a MTS 880 testing machine. Loading was increased in each step so that 1-5 unit of dial gage readings was obtained and recorded. The discrete force-indentation data were then plotted for analysis.

II - EXPERIMENTAL RESULTS AND DISCUSSION

Since there is permanent indentation at the contact point of composite shell during indentation test even for very small load, the unloading path is different from the loading path. In addition, experiment data show that the reloading path follows another path different from both loading and unloading, but it returns to the point where unloading starts. Thus, the whole contact processes will be depicted by three different stages.

317

2.1 Loading Stage

It was found that the Hertzian type 1.5 power law could be used in the loading stage for the indentation behavior of a ball on the composite laminated plate. The force F and indentation a is related by the following equation,

F = ka1.5 (1) where k is the loading regidity. For the case of composite thin shells, the above equation is also employed. Figure 1 shows that Eq. (1) could adequately fit the test data. The value k varies with different radii of curavture of the shell and steel ball. It means that the k is a function of curvaturs of the shell and indentation.

Table 1 summarizes the values for k, it shows that, in general, k increases as the radius of curvature of the shell increases for a given size of indentator. It is noted that for 12.7 mm (diameter) indentator on the [0/45/0/ - 45/0h. composite, the k value for Rt =39 cm shell is much smaller than the Rt =23 cm shell, as a matter of fact, it is the smallest of the four specimens with same stacking sequence. This phenomanon may be caused by the bad specimens which resulted from curing process. Another fact has to be pointed out here that for 6.35 mm ball, the k values are almost the same. It mean!! that for small size of indentator, the contact behavior for shells can be considered aR the same to that for a plate. When the radius of the indent at or increases, the curvature effect on the value of k become conspicuous.

2.2 Unloading stage

As in the indentation test for laminated plate [5], the unloading curve is different from the loading one significantly, and permanent deformation exists. Crook's power law

F = F ( a - a o )q m am - a o (2)

is also used here to model the unloading path. In the above equation, q is cho­sen to be 2.5, Fm and am are the contact force and corresponding indentation, respectively, when the unloading starts and a o is the permanent indentation. Equation (2) could also be written as

F = S(a - a o )2.5

where S = Fm/(am - a o)2.5

S is recognized as unloading rigidity. Reference [5] suggests that aCT is defined by fixing S as follows:

~ _ k ~CT - S

The permanent indentation ao could be obtained by

ao/am = 1 - (a cT /am)2/5 for am 2" aCT

a o = 0 for am < aCT

(3)

(4)

(5)

aCT is a material constants and is calculated by "area fit" of the unloading curve. Once aCT is determined, the unloading law follows. In the shell indentation, similar procedures lead to the value of aCT = 0.038mm for the material used. A sample unloading curve with test data is shown in Fig. 2. Good agreement is .... J.,,"" ... , ...... h.c.+UTo.on nrot-lirt.lPrl rnTVP ;;tntl t.p~t. TPRll1t..

318

2.3 Reloading Stage

Experimental data show that the reloading curve returns to the point where unloading starts, i.e., (Fm' am). Again, the reloading curve will be modeled by the equation similar to the plate indentation in [51. Hence during reloading stage

F=k1(a-ao)1.5 (6) where kl = Fm/(am - a o)1.5

Equation (6) fits the experimental results very well as shown in Fig. 3.

III - MODELING OF THE LOADING STAGE

In the preceding section, we know that the k and the critical indentation a cr

play key roles in the contact laws. The a cr is considered as a material property and is determined from experiment. If we can find an equation to predict the loading path, then the costly experiments could be reduced. A modified Hertz's contact law [91 is proposed here,

F = k(Ra, Rt , Ea, .. . )a1.5 (7) where k is a function of material properties as well as the radii of curvature

of the steel ball and the composite shell. It could be expressed as

k- C ~ - I-v2 + ...!... V Ji;.tii;

E. E2

(8)

where Es , Vs are Young's modulus and Poisson's ratio of the steel ball, E2 is the transverse modulus of the lamina, and RB, Rt are the radii of curvature of the steel ball and target shell, respectively. In Eq. (8), C is a constant which is determined from experiments. The values of C for the two different types of laminated thin shell are listed in Table 2. It can be seen that C is independent of the stacking sequence. The value of C is greater than 4/3 which is used in the Hertz law for two solid isotropic sphere. We may conclude that C takes account of the anisotropy and hollow property of the target shells. Sample loading curves according to Eqs. (7) and (8) is plotted in Fig. 4. Good agreements with the experimental data are demonstrated.

CONCLUSION

Contact laws of a solid steel ball and composite thin shell!; are derived through static indentation test. Hertzian law for the contact of two solid isotropic spheres is modified to simulate the loading stage of an indentor and composite thin shells. Good agreements show that the modified Hertzian law is adequate to describe the contact behavior of the material used in this research. It is suggested that further tests for other materials, e.g., glass/epoxy, are necessary to validate Eq.(8).

The contact laws derived herein is ready to incorporated with the shell finite element formulation to analyze the dynamic responses of a laminated composite thin shell subjected to impact.

319

ACKNOWLEDGEMENT

This work was supported by National Sci('nce Council and Foundation of Defense Industry Development under contract CS 77-021O-D006-19, Dr. D. G. Huang is the project monitor. The authors are grateful to Dr. C.T. Sun of Purdue University for his valuable suggestion in this research.

REFERENCES

1. Foreign Object Impact Damage to Composites, ASTM STP 568, (1973). 2. Dobyns, A.L., AIAA Journal, 19 5(1981) 642-650. 3. Shivakumar, K.N., Elker, W. and Illg, W., AIAA Journal, 23 3 (1985) 442-449. 4. Ramkumar, R.L. and Thakar, Y.R., J. of Engrg. Materials and Technology,

ASME, 109 (1987). 5. Yang, S.H. and Sun, C.T., ASTM STP 787 (1982) 425-449. 6. Tan, T.M. and Sun, C.T., J. of Applied Mech, ASME, 107 (1985) 6-12. 7. Chen, J.K. and Sun. C.T., Composite Structures, 4 (1985) 59-73. 8. Lee, L.J. and Huang, S.T., Proc. of International Congress on Composite

Materials, (1988) 839-846, Milan, Italy. 9. Goldsmith, W., Impact, Edward Arnold Ltd., London, 1960

Radius Type k(lO"' Nt/mm1. 5 )

of Steel of Stacking Radius of Target Shells Rt (cm) Ball Rs(mm) Sequences 10 15 23 39

6.35 A' 2.60 2.81 2.80 2.81 B' 2.64 2.71 2.78 2.75

12.7 A' 4.36 4.41 5.04 4.31 B' 4.47 4.79 5.39 5.01

2.54 A' 4.73 4.67 5.88 6.53 B' 4.86 5.99 6.25 6.21

* A' [n/45/0/-45/0h. B: [90/45/90/-45/90h.

Table 1 Loading Rigidity k for Different Cases

[0/45/0/-45/0hs [90/45/90/-45/90bs

6.35 1.70 1.68 12.7 2.12 2.30 25.4 2.01 2.16

Table 2 Values for Constant C in Eq. (8)

Hill

1m

EiIII

z \L GJ

200

a (mm)

Fig. 1 Sample Loading Curve for R = 25.4 mm

s

[0/45/0/- 45/0J 2s

R = 6 . 35 mm s

Rt = JOcm

CX(mm)

Fig. 3 Sample Reloading Curve

a (mm)

Fig. 2 Sample Unloading Curve

Hill [90/45/90/- 45/90] 25

:~ Rt = JOcm

J-R=J2 . 7mm s

2 - R = 6.35mm 5

400

200

cx(mm)

Fig. 4 Modeled Loading Path and Test Results

ABSTRACT

LARGE DEFLECTION INITIAL FAILURE OF LAMINATED RECTANGULAR PLATES

G.TURVEY,M.OSMAN

University of Lancaster - Department of Engineering Bailrigg - LA 1 4 YR LANCASTER - England

A f ini te-di fference version of the Dynami c Relaxat ion (DR) method is used to generate elastic solutions of the large deflection laminated Mindli n plate equati ons. These solutions are combined wi th the Tsai-Hill failure criterion to produce initial failure data for simply supported and clamped, specially orthotropic and cross-ply laminated square plates subjected to uniform pressure loading. Pressure and deflection versus (slenderness)' and span: thickness ratio plots are presented in order to quantify the effects of: plate thickness, in-plane edge restraint and membrane action on the initial failure response.

I-INTRODUCTION

Flat rectangular laminated plates may be regarded as the 'building blocks' of fibre-composi te structures. A complete understanding of the behaviour of flat laminated plates subjected to a wide variety of loading configurations is, therefore, essential for the development of efficient procedures for the design of fibre-composite thin-walled structures.

Although understanding and knowledge of laminated plate response has been developing steadily over the past two decades, it is, neverthe­less, still in its infancy when compared with the present situation for metallic plate response.

The response of laminated plates to transverse pressure loading is important from the practical standpoint - it represents the primary design load for many fibre-composite structures and components. How­ever, little is known about the failure of fibre-composite plates subjected to this type of loading. The reason for this is simply that

322

laminate flexural failure models are not well developed and remain to be validated by flexural failure tests on fi bre-composi te plates. Progress towards closing these gaps in our understanding of flexural failure response may be anticipated in the future, but the design of fibre-composite structures cannot await such developments, instead it must proceed, albeit with more caution, making use of the present state of knowledge.

As a preliminary to the development of collapse load data for pressure loaded laminated plates, studies of their initial failure behaviour, which may be considered analogous to first yield in metallic plates, are necessary and may be expected to provide data of value for preliminary design analysis.

This paper is concerned with the presentation of data derived from an initial failure study of pressure loaded specially orthotropic and cross-ply laminated square plates. The study represent s an extens ion of earlier work on thin plate failure analysi s by Turvey /1,2,3/. Small and large deflection solutions of the Mindli n laminated plate equations for the case of uniform pressure loading are combined with the Tsai-Hill failure criterion to generate initial flexural failure data - pressures and maximum deflections - for simply supported and clamped square plates. The results of the study serve to illustrate and quantify the effects of membrane action, plate slenderness and edge support conditions on the initial failure response and point the way to the rapid generation of approximate design data.

II - PLATE GEOMETRY AND MATERIAL PROPERTIES

Only square plates are analysed. Furthermore, they are considered to be either specially orthotropic or cross-ply laminated. Each plate is assumed to be layed-up from uni-directional CFRP pre-preg. The specially orthotropic plates are made from a single lamina. Table 1 gives details of the lamina elastic and strength properties.

III-LAMINATED MINDLIN PLATE EQUATIONS

The elastic solution of the laminated Mindlin plate equations constitutes a major part of the initial failure analysis. A brief summary of the large deflection versions of the equilibrium etc equations is given below. The small deflection versions are obtained by deleting the underlined terms.

3.1 Equilibrium Equations

N· + N' = 0 ; N" + N' 0 x xy xy y

Q. + Q' + N w ..

+ 2N w· , + N wit + q 0 ( 1) x y x xy Y

M" + M' - Q 0 M" + M' - Q 0 x xy x xy y y

in which the dot and prime superscripts represent respectively partial differentiation with respect to x and y. The other terms in Eqs.(l)

323

assume their usual meanings (see Reddy /4/ for further details including the derivation of Eqs.(1)).

3.2 Strain/CUrvature Equations

eO = u + ~(w. ) 2 eO = v' + ~ x y

eO = u ' + V + W w' eO 0 + w eO o + w' (2) xy xz x yz y

kO = 0' kO = 0' ; kO 0' + 0' x x y y xy y x

in which the superscript zero denotes strains and curvatures of the plate mid-plane.

3.3 Constitutive Equations

The constitutive equations for cross-ply laminates are:-

0

rr B12 0 r: [NX T' A12 ex + B11

0 0

0 kO Ny A12 A22 e y B12 B22 Y

N 0 0 A66 e:y 0 0 B22 kO xy xy

0

l[ °12 0

J lkO

["X 1'[B11

B12 ex + 011 x

My B12 B22 o e~ 012 °22 o k O (3 )

066 k:y M 0 0 B66 exy 0 0 xy

[::] . [:44 :,J ~~:l in which A .. , B.. and 0 .. are the in-plane/transverse shear, coupling and flexura~ rig\~ities iJspectively and are evaluated in the standard manner (see Jones /5/). The constitutive equations for specially orthotropic plates are derived from Eqs.(3) by setting the coupling stiffnesses, B .. , to zero.

1)

IV-TSAI-HILL FAILURE CRITERION

A number of failure criteria have been suggested for use with fibre-composite materials. A recent study by Turvey and Osman /6/ indicates that several of these criteria give broadly similar results for initial failure in flexure. Hence, the Tsai-Hill criterion is the only one considered here (similar results have also been computed for the Maximum Stress, Maximum Strain, Hoffman and Tsai-Wu failure criteria, but they are omitted for the sake of brevity). It may be expressed as follows:

F 0 2 + 2F 0 0 + F 0 2 + F 0 2 + F 0 2 + F 0 2 (4) 11 L 12 L T 22 T 44 LZ 55 TZ 66 LT

324

-2 -2 -2 -2 in which F = X F = -~X F = Y ; F44 = R

_211 t,c _212 t,c 22 t,c F = Sand F = T and the subscripts t and c attached to the

55 66 . f'" h strengths parallel and perpendJcular to the lbre-dlrectlon denote t e tensile and compressive values respectively. Similarly, the subscripts attached to the stresses in Eq. (4) refer to the principal material axes.

V-SOLUTION OF THE PLATE EQUATIONS AND SATISFACTION OF THE FAILURE CRITERION

Exact solutions of the small deflection Mindlin plate equations, i.e. Eqs.(1)-(3) with the underlined terms deleted, are computed using the Fourier series analysis outlined by Reddy /4/. Approximate small and large deflection solutions are computed using a finite-difference implementation of the Dynamic Relaxation (DR) method (see Turvey and Osman /6/ for a full description). The exact solutions are only valid for SS3 simply supported boundary conditions (see Fig.1(a)) and prim­arily serve to validate the small deflection DR analysis. The majority of results have been obtained for SS1 and SS2 simply supported and CC1

and CC2 clamped edge conditions (see Figs.1(a) and 1(b) respectively). They correspond to extreme combinations of flexural and in-plane rest­raint.

The particular elastic solution of the plate equations which just satisfies the Tsai-Hill failure criterion at one point (or occasion­ally several symmetric points) in the plate is determined by means of a simple scaling procedure for small deflection initial failure and by an interative re-analysis procedure for large deflection failure. Details of these two approaches are set out in Turvey and Osman /6/.

VI-NUMERICAL RESULTS AND DISCUSSION

In order to establish the accuracy of the approximate DR initial failure analysis, results obtained with the DR program were compared with corresponding exact results obtained from the Fourier series analy­sis. The results comparison is shown in Table 2 for thin and moderately thick plates. Good agreement for failure pressures, associated centre deflections and failure locations exists in every case. It is of interest to observe that the cross-ply plates fail at their centres whereas the specially orthotropic plates fail at the corners.

The large deflection initial failure results for various lay-ups, span : thickness ratios and support conditions are listed in Table 3. In order to demonstrate the effects of membrane action (the distinguish­ing feature between small and large deflection behaviour) on the initial failure response it is convenient to compare corresponding small and large deflection data. Fig. 2 (a) shows such a comparison for the S51 edge condition. The dimensionless failure pressure is plotted against what may be regarded as the square of the dimensionless slenderness. Clearly, the small deflection response is linear and the large deflection response is mildly nonlinear as the plates become thinner. It is also evident that for specially orthotropic plates small deflection theory

325

leads to unconservative failure pressure estimates when the plates are thin. The inverse situation applies for cross-ply laminated plates irrespecti ve of whether the lay-up is symmetric or anti symmetric . A similar form of plot, but for the centre deflection at failure, is shown in Fig.2(b). For specially orthotropic plates small deflection theory significantly over-estimates the deflections that arise at fail­ure. However, for cross-ply lay-ups the small deflection analysis predicts failure deflections which are only marginally unconservative. Fig.2(b) also suggests that the maximum deflection at failure is very nearly proportional to the square of the plate slenderness for both small and large deflection analyses. This feature, if generally true, could be of significance for the production of simplified design data.

The effect of the degree of in-plane restraint on the initial failure pressures of simply supported plates is shown in Fig.2(c). Clearly, the failure pressure increases as the plate becomes thinner and as the in-plane restraint increases.

A similar plot to that of Fig.2(a), but for plates with CC1 clamped edges, is shown in Fig. 3 (a) . Again, the small deflection response is linear and the large deflection response is very nearly so, because the effect of the rotational edge restraint is to limit the development of significant membrane act jon prior to initial failure. The small deflection initial failure pressures for thin specially ortho­tropic plates are marginally unconservative compared with the large deflection values. For cross-ply plates, as was observed for 551 edge condi tions, the inverse situation applies. Ini tial failure pressure versus span : thickness ratio plots for plates with CC1 and CC2 edge conditions are shown in Fig. 3(b). It is obvious that the effect of increasing the degree of in-plane edge restraint is to raise the initial failure pressure, especially as the plates become thinner. However, the in-plane edge restraint effect is much less marked under clamped edge conditions (cf. Figs.2(c) and 3(b».

VII -CONCLUDING REMARKS

Initial failure data for uniformly loaded specially orthotropic and cross-ply laminated square plates with simply supported and clamped edges have been computed using small and large deflection Mindlin plate theory and the Tsai-Hill failure criterion.

It has been shown that failure pressure versus square of slender­ness plots are linear for small deflection analysis and could be used for I conservative I design analysis for cross-ply plates. Only when the plates are very thin does large deflection analysis provide less conservative failure pressure data, particularly for simply supported plates. Indeed, the linear nature of the failure pressure response might be exploited for the generation of approximate design data. For any cross-ply lay-up the failure pressure - square of slenderness response line could be defined from just two small deflection failure analyses - one, pay, for a moderately thick and another for a thin plate. The straight lj ne joining these two failure pressure data points would then represent the design curve for all plates. Maximum

326

deflection design data could be derived in a similar manner, but the data would be marginally unconservative.

VIII-ACKNOWLEDGEMENTS

The second author wishes to thank the Sudanese Mini stry of Hi gher Education for providing a Studentship to enable thi s research to be completed. Both authors wish to record their thanks to the Department of Engineering for supporting their work. Finally, thanks are due to Mrs. Kathryn Rucastle for typing the paper and t.o Mrs. Audrey Parker for preparing tracings of the figures.

IX-REFERENCES

1. G.J. Turvey, Int. J. Solids and Structures, 16 (1980) 451-463

2. G.J. Turvey, J. Composite Materials (Supplement), 14 (1980) 1-14

3. G.J. Turvey, in "ICCM3", (A.R. Bunsell, ed) (1980) 291-304, Paris

4. J.N. Reddy, "Energy and Variational Methods in Applied Mechanics", John Wiley & Sons, New York, (1984) 389-401

5. R.M. Jones, "Mechanics of Composite Materials", McGraw-Hi 11 Koga­kusha Ltd, Tokyo, (1975) 154-155

6. G.J. Turvey and M.Y. Osman in "ICCS5" (I.H. Marshall ed.) (1989) Paisley.

TABLE 1

Material Properties of Uni-Directional CFRP Pre-Preg

12.31

34.6

Note:

(a) Elastic Modular Ratios

0.526 0.526 0.314

(b) Strength Ratios

x /Y c t

38.7

Y /Y c t

1.000

245.7

1.543

0.24

1.984 1.984

327

TABLE 2

Comparison of DR and Exact Small Deflection Initial Failure Analysis Results for Uniformly Loaded Square Plates with SS3 Simply Supported Edge Conditions

4 4 w /h x/a y/a h /a Lay-Up qa /E h k s

T 0 c 0 0

0 5.669* 0.0819 0.5 0.3 3 0.10 (Ortho- 5.584** 0.0812 0.5 0.3 3 tropic) 154.98 1.8615 0.5 0.3 3 0.02

150.20 1.8109 0.5 0.2 3 A2 2.438 0.0614 0.5 0.5 2 3 0.10

(0°/90°) 2.379 0.0602 0.5 0.5 2 3 60.780 1.4104 0.5 0.5 2 3 0.02 59.409 1.3846 0.5 0.5 2 3

S4 5.478 0.0810 0.5 0.5 4 3 0.10 (0°/90°) 5.359 0.0799 0.5 0.5 4 3 90%°) 150.27 1.8419 0.5 0.5 4 3 0.02

147.58 1.8167 0.5 0.5 4 3 A4 4.477 0.0714 0.5 0.5 4 3 0.10

(0°/90°/ 4.367 0.0701 0.5 0.5 4 3 0°/90°) 111.58 1.5597 0.5 0.5 4 3 0.02

108.79 1.5274 0.5 0.5 4 3

* DR quarter-plate analysis (5 x 5 uniform non-interlacing mesh)

** Exact analysis (Series solution)

Note:

k = lamina number (starting from upper loaded surface) s = lamina failure surface (1 denotes upper and 3 lower surface) Shear correction factors = 5/6 Failure initiates either at a unique centre point or at two points of symmetry. In the latter case only one of the points is given in the table.

328

TABLE 3

DR Large Deflection Initi.al Failure Results for Uniformly Loaded, Square, Simply Supported and Clamped, Speci.ally Orthotropic and Cross-Ply Laminated Plates.

(a) SSl and SS2 Simply Supported Plates

a/h 4 4

w /h x/a y/a Lay-Up qa /E h k s 0 T 0 c 0

0 10 5.573* 0.0804 0.5 0.3 3 (Ortho- 5.514** 0.0783 0.5 0.3 3 tropic) 20 23.177 0.2886 0.5 0.3 3

23.509 0.2546 0.5 0.2 3 30 51.300 0.6062 0.5 0.3 3

56.770 0.4568 0.5 0.2 3 40 89.92 1.0024 0.5 0.2 3

111.32 0.6487 0.5 0.1 3 50 138.38 1.4319 0.5 0.2 3

189.61 0.8219 0.5 0.1 1 3 A2 10 2.419 0.0608 0.5 0.5 2 3

(0°/90°) 4.100 0.0652 0.4 0.5 2 3 20 9.800 0.2282 0.5 0.5 2 3

17.400 0.2078 0.4 0.0 1 30 23.270 0.5059 0.5 0.5 2 3

38.000 0.3953 0.5 0.0 1 40 48.604 0.9137 0.5 0.5 2 3

69.303 0.5819 0.5 0.0 1 1 50 99.967 1.4660 0.3 0.5 2 3

125.56 0.7695 0.5 0.0 1 S4 10 5.444 0.0804 0.5 0.5 4 3

(0°/90°/ 5.317 0.0774 0.5 0.5 4 3 90%°) 20 23.399 0.2957 0.5 0.5 4 3

24.000 0.2649 0.5 0.5 4 3 30 56.000 0.6528 0.5 0.5 4 3

69.811 0.5295 0.5 0.3 4 3 40 110.53 1.1365 0.5 0.3 4 3

172.51 0.8228 0.5 0.3 4 3 50 198.52 1.7074 0.5 0.3 4 3

379.51 1. 1351 0.5 0.3 4 3 A4 10 4.467 0.0712 0.5 0.5 4 3

(0°/90° 5.130 0.0740 0.5 0.5 4 3 0°/90°) 20 17.860 0.2544 0.5 0.5 4 3

21.599 0.2492 0.5 0.5 4 3 30 41.258 0.5531 0.5 0.5 4 3

63.446 0.5146 0.5 0.5 4 3 40 81.000 0.9759 0.5 0.5 4 3

168.30 0.8311 0.3 0.5 4 3 50 157.00 1.5604 0.5 0.5 4 3

370.20 1.1423 0.3 0.5 4 3

329

(b) CC1 and CC2 Clamped Plates

a/h 4 4

w /h Lay-Up qa /E h x/a y/a k s 0 T 0 c 0

0 10 4.114+ 0.0213 0.5 0.0 1 (Ortho- 4.139++ 0.0214 0.5 0.0 3 tropic) 20 21.000 0.0678 0.5 0.0 1

21.234 0.0687 0.5 0.0 3 30 50.670 0.1441 0 .. 5 0.0 1

51.597 0.1475 0.5 0.0 3 40 93.271 0.2483 0.5 0.0

94.000 0.2564 0.5 0.0 50 150.60 0.3745 0.5 0.0

150.00 0.3995 0.5 0.0 A2 10 2.918 0.0225 0.5 0.0 1

(0°/90 0 ) 2.938 0.0227 0.0 0.5 2 3 20 11.672 0.0709 0.5 0.0

11. 713 0.0713 0.0 0.5 2 3 30 26.864 0.1534 0.5 0.0 1 1

26.375 0.1524 0.0 0.5 2 3 40 50.160 0.2689 0.5 0.0 1 1

47.459 0.2687 0.0 0.5 2 3 50 86.734 0.4142 0.5 0.0

74.961 0.4165 0.0 0.5 2 3 54 10 5.224 0.0276 0.5 0.0 1 1

(0 0 /90 0 / 5.289 0.0280 0.5 0.0 4 3 90 0 /0 0 ) 20 25.798 0.0862 0.5 0.0

25.798 0.0863 0.5 0.0 4 3 30 62.000 0.1819 0.5 0.0

62.000 0.1834 0.5 0.0 4 3 40 116.39 0.3162 0.5 0.0 1 1

112.70 0.3173 0.5 0.0 4 3 50 196.75 0.4858 0.5 0.0 1

180.00 0.4921 0.5 0.0 4 3 A4 10 6.118 0.0318 0.5 0.0 1 1

(0°/90 0 / 6.118 0.0319 0.0 0.5 4 3 0 0 /90 0 ) 20 24.831 0.0899 0.5 0.0 1

25.154 0.0913 0.0 0.5 4 3 30 56.741 0.1871 0.5 0.0 1 1

56.393 0.1877 0.0 0.5 4 3 40 104.12 0.3214 0.5 0.0 1

101. 70 0.3273 0.0 0.5 4 3 50 174.36 0.4891 0.5 0.0 1

160.20 0.5039 0.0 0.5 4 3

* 551 results; ** 552 results; + CC1 results; ++ CC2 results.

Note: All DR results are computed with a 5 x 5 uniform non interlacing mesh over one quarter of the plate.

Failure initiates either at a unique centre point or at two points of symmetry. In the latter case only one of the two points is given in the table.

330

"Y r .

::!:D· Hx' ° ','0

-..x,u

·'°0 y=o

w=o

H '0 x • '0 y

NX'OO v =0

w= •

H,"

'1'0 I,W Nx,' N,' W "x'H,'O u' "y'W"X'H,"

ill ill

(a) Simply supported edges

II" 1l:,::v:,X=',=O [[1

NX'OO H ,. x, V =0

I '0 x . '. y

Nx,' N,' W "X"1" [ [2

(b) Clamped edges

Fi g. 1 - Plat_e edge condi tIon combInatIons

tI. -: lMG[ DEFl[[lION --: SHAll DEFt[C liON S'

110 D : SP[CIAllY OATHOTROPIC PLAIE

AI,A': ANIISYHHITR1C ~~I~~~~~:'

to S.: SYHHETRIC FOUA-lAYEA PlATE

~n [r tl• AI

U

II

I

• • (N~r • IT

(a) Pressure versus (slenderness)' for SS1 simply supported plates

~ ~.

H

I-I

1-0

... 04

H

0-2

I-I 0

- : lAAG[DEFlECIION ---- 'SHAll DEFl[(1I0N

o : SP[CIAllY O",NOTAOPIC PLAn 0

A2,A" ~ ANflSYHHETAIC l~~[~NBt:T~~· ~,,~ ~~ ~

S4 : SYHHURt( FOUR-unA PLATt: ,::J

" " ",'

,,:' ~

,,.,,,'"

"'4,0 .. AI,AI

(b) Associated centre deflection versus (slenderness)' for SS1 simply supported plates

200,------------------.7"""1

150

SO

-- : 551 EDGE CONDITIONS ---- I 552 EDGE CONDITIONS

o : SPinALLY ORTHOTROPIC PLAT! A2,AlI ANTlSYHHETRIC 1 liD AND FOUR­

lAYER PlAT£S 54 I 5YHHfTRIC fOUR-lAYER PLATE

,54 ' ...... /f

" 1/ 'I

/1 " " ,/

II /1

" " " 'I

1/

10 20 ~ h.

" " II I'

" I' 1/ 'I

,'1

JO 40 10

(c) Pressure versus span: thickness ratio (a/h ) for SS1 and SS2 simply sup~orted plates

Fig.2 - Initial failure data for simply supported plates

331

~r-----------------------~ - ,CC' lOG! CONDI1IIIIS - - - - : (( I [Oli[ (ONDITIONS

o 'SP[(IAlLY ORTHOTROP\( PlAT( AI.'" : ANTlSYM"ETRIC fWD AND FOUR­

LAYER PLATES 5" 'SVH"URI( fOUR-lAYf.RPUTt

12' r---==, -;-' .. =.I;-;.:;:";o";;'""n."';----------"..-, '" -_.-j SHALl. DUUt1l0N

o 1 SPE(lAllY ORTHOTROPIt PLAT(

110 A2,Att = ,,"tISY""'TRIC 1~~(\N~lr:~[Rs' 54 I SVHHURI( FoUR·LAYER PLATE

..

.. 10

'~~L-__ ~ __ ~ __ -L __ -L __ -L __ -J

I J •

(!.ffi,V (a) Pressure versus (slenderness)'

for CC1 clamped plates

"" q.

(r~

"

'L-_~~ __ ~ ____ ~_~~_~~ o 10 2. 0.]1 4. SO

h.

(b) Pressure versus span : thick­ness ratio (alh ) for CC1 and CC2 clamped p.la~es

Fig. 3 - Initial failure data for clamped pLates

POST·BUCKLING OF FLAT STIFFENED GRAPHITEIEPOXY PANELS UNDER CYCLIC COMPRESSION

Y. FROSTIG, A. SEGAL',1. SHEINMAN, T. WELLER

Faculty of Civil Engineering, Structures Division Technion, 1.1. T. - 32000 HAIFA - Israel

• Israel Aircraft Industries - BEN GURION - Israel

Results of expeI'lIrental and analytlcal study of tile post-buckl­Ing of selected flat stIffened Graptute/EPoxy panels, loaded in axIal cO!Ji)resslon are presented. TIle lIll tlal bucklIng, POSt-l:lUCkl ing and failure characterlstlc are deflned and descrIbed nle axIal load pattern Included statIC test to faIlw'e and CYClIC loadlng up to a total of 250,000 cycles. TIle analytIcal part COnsIsts of results from NASTRAN - a flmte eleIrents package, PB:XMP an In-house developed program wtuch performs lInear and non-lInear buckling analysis and apprOxllmte Irethods.

A reductlon In the 1m tlal bucklIng load after cyc ling with an Increase In tile faIlure load was observed The study proves tllat the performance of such structures IS only slightly affected by CYClIC CO!Ji)reSslon.

MetallIC stIffened plates and shells are designed to buckle below deSIgn ultImate load. Vast Information, experiIrental and analy­tIcal, exists on tile post-buckling behaVIor of Iretallic panels howev­er, only limited data IS aVailable on tile behaVior of laminated compoSite panels and tilelr abIlIty to witllstand loads In the post­bucklIng range. As a resul t, composIte structures are currentl y deSIgned not to buckle below design ultImate load

A large number of publicatIOns dealing WIth post-buckling behaVi­or of Iretallic panels under corrpression and shear loads appears In tile literature, see for exauple [1) and [2). lXlring tile last few decades the classical buckling problem was extended, to SOIre extent to anisotropic plates, see Ashton, [3), Agarwal [4J on shear panels and Starnes [5). Leissa [6) Cites a few references, on stIffened panels wluch use analytIcal procedures to predict tile post buckling behaVIor. Studies on fatigue effects are rarely found in tile open Ii terature and aroong the few we can cited are refs. [7-8).

334

The maln objectives of the study werej evaluation of the post­buckling potential of stiffened GRIEP panels under compressive loads, identification of the failure modes, fatigue effects studies and development of analytical prediction tools of the initial buckllng and failure modes and to establish the efficiency of the various stiffener shapes.

The marruscript presents the testing program, experimental resul ts for the J-type panels, analytical predictions, conclusions and recom­mendations. The experimental results for the I-type panels appear in a previous paper, see [8J.

Seven panels tested, with I and J shape stiffeners, see fig. 1, were designed and manufactured at Israel Aircraft Industry and tested at tt~ Aeronautical Engineering Structures Laboratory of the TeChn­ion, Haifa The specimens were fabricated with AS4/3502, graphite epoxy tape material Which is a 350<>f' cure type. The dimensions and the stacking sequences of the stiffeners appear in fig. 2.

The manufacturing process consisted of one curing cycle with the stiffeners integrally laid into the skin and cocured with it. Follow­ing cure the specimens were ultrasonically inspected The defects were mostly found in the cap-web joints and at the interface between the stiffener and the skin

The ends of each specimen were potted in an epoxy-resin material mixed with glass fibre, encased in a U shaped aluminum channel. The loaded ends were machined flat and parallel to permit uniform axial compressive loading. The panels were painted White on both sides to enhance visibility of brooming of the black Graphite fibres and the molre-fringes of tIle out of plane displacement pattern

Strains, displacements and skin bUckling IOOde shapes were monit­ored with the aid of back to back bonded strain-gages, displacement meters and the moire-fringe technique. The cyclic loading was applied between 0 to Pmax at a frequency of 2 Hz in blocks of constant anpli­tude corrpression and with a maxllIl.1lIl load ranging from twice the ini tial bUckllng load up to a maxilIl.llIl of 70'1. of the fall ure load of the reference panel. The experimental results for the J-type panels are described next.

'lEST RESULTS

J - PANELS

The first speCimen, in this group, panel J4, was statically loaded to failure and served as a reference. The Initial buckling load was determined USlng the procedure described in ref. [9), and the mode pattern consisted of an unsyrJIJetrical buckling mode shapes, 6-5-5 halfwaves in the various bays. The collapse pattern consisted of cracks and fibres breakage at the defects locations and 450 cracks in the skin Manufacturing defect dimensions remain unchanged during the testing program.

Panel J6 was the first f)ne successfully tested under the cycllng

335

conpression loading. '!be panel had SCIre marru.facturing defects which did not affect the behavior or the failure pattern. '!be dimensions of the defects remained lmChanged throughout the cyclic loading. '!be only damage, as a result of the cyclic loading, was in the fonn of broouung of fibres at the web-flange (SKin) intersection point at one of the side stiffener. '!be initial ruckling load level was, nearly twice as high as that of the reference specinen, J4. It had consisted of a s}'lJJDetric ruckling mde sbape pattern with five halfwaves at each bay. In the post-Wckling range the averaae strains were redUced with an increase in the IlUIItIer of cycles while the bending strains were unaffected. Failure load was about 10:1. higher tilan that of the reference specimen Collapse pattern which was initiated at the damaged side stiffener, was similar to that of J4.

'!be last panel in the program, subjected to the cyclic compress­ion loading, was panel J7, which was identical with the previous ones rut defects free. Initial Wckling load was about 20:1. higher tilan the reference specimen and consisted of an Ul'lS}'IJJDetrical pattern, 5-6-6 halfWaves in the various bays, see fig. 3. '!be Ul'lS'}'IJJJIetrical pattern was probably due to the nonunifonnity in the longitUdinal boundary conditions. '!be panel was cycled to 250k cycles without any observed damage, at load levels deep in the post-Wckling range. 'nle initial ruckling load level decreased as the IlUIItIer of cycles was increased. However, the failure load was about 25:1. higher tilan that of the uncycled panel. '!be collapse occurred in the three stiffeners sinnl t­eneously, in the fonn of delam1nations in the caps and mtrix cracks wi th fibres lreakage in the skin at mid-span.

'nle stiffeners behavior followed the trends of an Ul'lS}'IJJDetrical column ruckling mde since the section as a Whole was partially fixed in the out of plane direction while its cap was fully fixed in the in plane direction, see fig. 4.

'nle skin average strains at the center of the side and mid bays were affected by the cycling compression, see fig. 5. 'nle load versus the skin bending strains, at the various bays before and after cycli­ng, followed the knoWn trends of post-l:Uckling :behavior of netallic panels with ifiIlerfections, see fig. 6. 'Ibe ruckling mde shape patte­rn in the central bay charlged frOm five to six halfwaves, during the strength residual test, causing a shift in the curve, see fig. 6, a known Iilenonenon in the non-linear :behavior of plates, see [10].

'Ibe load-bending strains curves of the various stiffeners were nearly identical, neaning that the loading was s}'lJJDetrically distrib­uted, see fig. 7. 'nle. stiffener :behavior followed the trends of classical ruc~ing of colums with inIlerfections. 'nle average strain distrirutions at the side and central bays before and after cycling appear in fig. 8. Strain concentration in the vicinity of the stiffe­ners was observed in the post-ruckling range. 'nle compression strains, in the center of the side bay, deep in the post ruckl ing range, were even changed to tensile ones.

The concentration of strains in the vicinity of the stiffener lead to the use of the "effective Width" concept. Side and central stiffeners "effective Width" were not the sane, see fig. 9, and cycling had lowered the values.

336

AHALYTlCAL RliSULTS

The analytical investigation included lnitial buckling analysis predictions and IOOde shapes using MSC NASTRAN [111, Pa:Xl'IP (Post -Buckling of cx:Mposite Panels) an in-house developed code [12, 131 and an approximated nethod based on an isolated bay with rotational springs buckling analysis.

The analytical roodel used the 12 and J7 data, and assumed clarIll­ed longitudinal ends and non-uniform sKin thiclmess in the bays, ln all nethods. The stiffeners, in NASI'RAN and pBXIfF' were considered as part of the roodel, while in the approximate nethod as concentrated rotational springs whiCh were located at the web sKin intersection pomt. A uniform load distrirution was considered in all nethods. A uniform end shortening loading pattern was checked using pBXIfF' and the approximate nethod The resul ts are presented in tabl e 1.

The fallure loads were predicted assuming ruckling of a sinply -supported colurm. The equivalent properties were calculated consider­ing an "effective width" of 90 nm of the attached sKin as well as the stiffener section The Simply-supported end conditiOns were used since tt~ U-Channel, at the ends, see fig. 1, was not capable of preventing the stiffeners rotations at the ends. The predicted loads were about 151.-201. higher than the experinental ones using the princ­lpal roonents of inertla

An experinental and analytical investlgation was conducted to study the ruckling and post-buckling behavior of GRIEP stiffened panels loaded in cyclic compression I and J type stiffeners were considered The conclUSiOns, wi ttl respect to the objectives listed in the introduction, are as follows: (1) The panels used in the program exhiblted post-buckling capability and behave similar to netallic panels. The design consisting of thin sKin and a heavy stiffener proved to be efficient since it had allow­ed the sKin to enter deep into the non-linear range. (2) The initial ruckling load of the sKin was affected by the cycling and had decreased as the number of cycles increased However the failure loads after cycling were even increased conpared with the uncycled specinen (3) The I-shaped stiffener panel configuration proved to be IOOre efficient compared with the J-shaped with the respect to strength­weight ratlo. The J-stiffener weight was about 711. of the I-section, while its failure loads were about 521.-571. of the I-type ones. A J­type panel total weight was about 861. of the I-type one. The 1-sti ffened pane 1 s exhibited higher ruckling strength and small deformatlons compared with the J configuration (4) The initial buckling predlctions, assuming a uniform load distri­bution, were satisfactory in all nethods, excluding NASTRAN results for the I type panels. Predictions using the end-shortening load pattern yielded hlgher values compared with the experinental results. The analytical buckling roode shapes pattern was very similar to that of the experinental one. (5) Failure roode, for both types of panels, conslsted of colurm

337

bucKlmg IOOde W1 th lllperfect10ns. No separatlOn of the sK1n from the stiffener was observed and no crippl1ng 1n the st1ffeners was detect­ed. (6) Pred.lctions of tile fa1lure loads USlng an eqtllvalent sectlOn, which consisted on the stiffener sectlon and an attached "effectlve width" sKin of O. 35b, proved to be satisfactory. (7) '!he failure pattern consisted of collapse of three stiffeners out of the four and was usually lnitlated at one of the slde stlffener and proceeded to the center ones. (8) '!he cycllc loading affected the membrane straln d.lstrlbutlons 1n the various bays. '!he strains at the center of the bays were reduced while those in the vicinlty of the stiffener were lncreased as the number of cycles was increased. '!he reduction in the strallls at the center of the bays, 1n some cases, changed the compress10n to tens10n due to excessive deformations. (9) Manufacturing defects dimensions remained unchanged under the cycling conpression loading. Failure was usually 1nit1ated at these locations but the overall behav10r was unaffected.

1. Wagner, "Ebene blechwandtrager sehr d:uIlmen stegblech", Flugtechmk und IOOtorluftschiflahrt, Ed. 20 1929.

2. Koiter, W., T., "On the stabll1ty of elast1c eqtl1llbrlurn" (In Dutch), thesis, Delft H. J. Paris, Amsterdam, 1945. English trans­lation, Air-force Flight Dyn. Lab. Tech Rep. AFFDL-TR-70-25, WPAFB, 1976.

3. Ashton, J., E., and Waddoups, M., E., "AnalYSlS of anlsotrop1c plates". Journal Of Conposite Materlals. 3(January 1969) 148-165.

4. Agarwal, B., L., "Post-buckling behavlor of conpos1te shear webs". AIAA Journal, Vol. 19, No.7, pp. 933-939, 1981.

5. Starnes, J., H., Jr and Rouse, M., "Post-buckllng behav10r of selected flat stiffened GraItute-Epoxy panels loaded 1n corrpress1-on" . Proceeding of the AlAA/ASME/ASCE/AHS 23rd SOO Conference. Lak.e 'rahoe, Nevada. May 1982.

6. Leissa, A., W., "Analysis of lanunated conposite plates and shell panels" , AF'NAL-TR-85-3069, AF Wrlght Patterson Aeronautical lab. Jan. 1985.

7. Weller, T., Kollet, M., Llba1i, A., and Singer, J. "DuraJ:.nl1ty under repeated bucKling of stiffened shear panels", 14th lCAS Conference, Toulouse, France, 932-942, September 19tY*.

8. Segal, A., Siton, G., Weller, T., " Durability of GralDite-Epoxy stiffened panels under cycling post-buckling conpreSS1on load.lng" Froc. ICa1 VI/EX:X-""M, London, July 87.

9. Sonsa, M., A., For, W., C., Walker, A., C., "Review of experlln:m­tal techniqtle for thin-walled structure liable to bucKllng part I and II " EKperlmental techniqtles, Sep. and Oct. 83.

10. Supple, W., J., "Structural Instab1l1ty", (Jluvers1ty of SUrrey, Ipc Science and Technology Press, Sep. 1973.

11. Macneal R., H., "MSC/NASI'RAN." Macneal-Schwendler Corp., Los Angeles, CA., 1983.

338

12. Sheiman, 1., F'rostig, Y., "Post-buckling analysis of stiffened laminated panels", Journal of Applied Mechanics, ASHE, Vol. 55, Sep. 1988, pp. 635-6l:lO.

13. Sheinmn, 1., Frostig, Y., Segal, A., ""Bifurcation Buckling Analysis of Stiffened Composite ~nated panels", in Studies in Appl ied mechanics 19 - "Buckl ing of Structures, 'Iheory and Exper­iments", the Josef Singer Anniversary Volume, Elsevier, Amsterdam, Oct. 1988, pp. 355-380.

Panel 12 Panel J7 Code /

Method End-sbOrt Uniform :End-sbOrt Uniform

Nastrar. 1.308 1. 2M Pbconp 1.%2 1. 036 1. 521 1. 275 Approx. 1. 530 1.0~ 1. 618 1. 309

Table 1: Predicted initial buckling load with respect to the experimental ones, for panels 12 and J7, using the various codes with uniform and uniform end-sbOrtening load distributions.

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339

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340

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FINITE ELEMENT ANALYSIS OF COMPOSITE PANEL FLUTTER

L.C. SHIAU, D.H. TSAY, L.J. LEE

Institute of Aeronautics and Astronautics National Cheng Kung University - 70101 TAIWAN - Republic of China

ABSTRACT

A finite element approach to panel flutter problem of finite laminated plate has been developed. Linear small deflection structural theory and quasi-steady aerodynamic theory are employed for the analysis. In order to study the effect of panel configurations on the flutter behavior of a composite laminated panel, a fully compatible general quadrilateral plate bending element is used. The effects of fiber orientations, stacking sequences, anisotropic properties, boundary conditions, aspect ratios, and panel configurations are studied. Calculations show that the flutter stability can be improved if the fiber is properly orientated.

INTRODUCTION

Panel flutter has been known as one of the most important problems in the design of aircrafts, missiles, launched vehicles, and spacecrafts. Extensive progress in flutter tests and theoretical studied has been achieved in the past thirty years (1,2). Recently, because of their superior strength-to-weight and stiffness-to-weight ratios, as compared with the conventional materials, composite materials has been widely used in the aeronautical industries for replacing aluminum alloy in the aircraft structures for the purpose of weight saving. At current stage, in the high performance aircraft and spacecraft, composite materials are mostly used on the skin of wings and fuselage of an aircraft. Since the skin plates fabricated from composite materials consist of individual lamina bonded together and the principal direction may differ for each lamina, the flutter behavior of the composite panel may be more complex than an conventional panel. Some investigators [3,4] have studied the effects of filament angle and direction of airflow on the flutter behavior and their work were concentrated on a particular stacking sequence or boundary condition which may not suit for the practical use.

342

In this study, a finite element approach to panel flutter problem of finite laminated plate has been studied. Linear small deflection structural theory and quasi-steady (linear) aerodynamic theory are employed for the analysis. In order to study the effect of panel configurations on the flutter behavior of a composite laminated panel, a fully compatible general quadrilateral plate bending element assembled from four partially constrained linear curvature compatible triangles [5] is used.

I - FORMULATION AND SOLUTIONS

Consider a thin flat plate of length a, width b, thickness h, and density p, with supersonic air flow passing over the upper surface as shown in Fig. 1. The differential equation of motion for anisofropic plate is

(1)

where D~js are the flexural and torsional rigidities of the plate and p is the aero­dynamic pressure. For sufficiently high supersonic speeds (M > 1.6), the aerody­namic pressure can be described by the quasi-steady aerodynamic theory:

p(x,y, t) = (2)

where U is the air flow velocity, qo = -!PaU2 is the dynamics pressure, Pa is the air density, and M is the Mach number.

Assuming the displacements are exponential functions of time

w = W(x,y)e vt

and introducing nondimensional variables

x x= -,

a

(3)

(4)

where v is a complex number and ET is the transverse (smaller) elastic modulus of the composite plate, Eqs. (1) and (2) become

a4w a4w a4w a4w dll aX4 + 2d12 aX4ay2 + d22 ay4 + 4d16 aX3ay

a4w a4w pa4v 2 M2 - 2 + 4d26 aXay3 + 4d66 aX2ay2 + ETh2 (1 + j.LC M2 _ I)W (5)

pa4v 2 aw +--j.LC "X =0 ET h2 u

where u c=-va

343

,

In the finite element method, the stiffness equations of motion for a plate for Eq. (1) under the influence of elastic, inertial, and aerodynamic forces may be written as

[K] {q} + [M] {ij} + [A] {q} = {O} (6)

where [K], [M], and [A] are, respectively, the stiffness, mass, and aerodynamic matrices assembled for the whole finite element system. The vector {q} contains the modal degrees of freesom for the whole panel system. After separating the aerodynamic matrix into aerodynamic force matrix [A J] and aerodynamic damping matrix [Ad] and with some rearrangement, the finite element form of Eq. (5) becomes

([K] + A [AJ]- k [M]) {q} = {O} (7)

in which

(8)

and 42M2 2 k = _pa v _ A __ -_(a)

E T h2 M2 - 1 U (9)

are the non dimensional dynamic pressure parameter and eigenvalue, respectively.

Eq. (7) represents an eigenvalue problem. For zero flow velocity, A = 0, the eigenvalues, k, are real. As the flow velocity increases from zero, two eigenvalues will usually approach each other and coalesce to kCT at a value of A = ACT> which is a critical value of dynamic pressure, and become complex conjugate pairs for A > ACT. Once A > ACT' the panel motion becomes unstable.

In general, due to the unsymmetric nature of the aerodynamic force matrix [AJ]' it is relatively time consuming to solve the flutter problem for a large system of equations as the case in this study. Hence, the modal method proposed by Rossettos and Tong [ ] is employed to reduce the order of the eigensystem to a rather small size.

II. RESULTS AND DISSCUSION

Fluttter solutions are obtained for symmetric laminated plates with material properties typical of graphite-epoxy (~ = 14, G LT = 0.53ET, VLT = 0.3)

and boron-epoxy (~ = 10, G LT = 0.333ET , VLT = 0.3).

As mentioned before, when the dynamic pressure increases from zero, two natural modes approach each other and coalesce at a value of ACT at which the flutter begins. Fig.2 shows the difference of the coalescent characteristic between isotropic and composite square panels. Due to the effect of directional stiffness of

344

the composites, the occurence of the flutter of the composite panel ill due to the coalescence of the 1st and 4th modes which is in contrast to the usual case of the isotropic panels where the coalescent modes are the lot and 2nd • The coalescent modes may differ from case to case for composite panels with different geometric and material properties such as ELI ET ratio, fiber orientation, flow angle, etc.

Flutter boundaries are shown as a function of fiber orientation in Fig. 3 with different number of layers. The highest flutter boundary is obtained with fibers aligned with the x-axis; rotating the fibers away from the x-axis results in a continuous reduction in flutter boundary for values of (J up to 90°. Since the bending-twisting stiffness terms have a pronounced destabilizing effect on the flutter boundaries and these coupling terms become smaller as the number of layers increases, the flutter boundaries can be improved with a larger number of layers as observed in the figure. Fig. 3 also shows the effect of boundary condition on the flutter boundaries. As expected, the fully clamped panel has the highest flutter boundaries due to its stiffness increased by the fixed edges

Fig. 4 shows the effect of the cross-flow angle on the flutter boundaries. At (J = 0°, cross-flow has significant destabilizing effect on the flutter boundary. This is due to the fact that the plate stiffness is gradually reduced in the flow direction as the flow angle increases. When cross-flow angles increase form 0° to 20° and 40°, the peak values of the flutter boundary are obtained at (J = 30° and 45°, respectively. Therefore, with larger cross-flow angles, the fibers should be orientated at an angle slightly higher than the cross-flow angle in order to increase the bending stiffness in that direction.

The effects of elastic modulus ratio on the flutter boundary for different fiber angles are depicted in Fig. 5. Increasing the modulus ratio increases the natural frequencies of the plate which in turn rise the flutter boundary, However, because the component of ELI ET in the flow direction becomes smaller as the fiber rotating away from the x-axis, the rate of increase of the flutter boundary is reduced with increasing fiber angles.

In real structures, panel configuration may not always be rectangular. Fig. 6 shows the effect of panel configuration on the flutter boundaries. The results show that panel with 8 = 110° has the highest flatter stability.

CONCLUSIONS

The finite element method has been used to study the flutter behavior of composite laminated panels. Because of its versatile applicability, effects of com­plex panel configuration, anisotropic material properties, uneven thickness, various boundary conditions, and flow angularity can be easily included. The parameters studied include the number of layers, fiber orientations, cross-flow angles, and panel configurations.

ACKNOWLEDGEMENTS

The financial support by National Science Council through contract No. CS77-0210-0006-20 ill gratefully acknowledged.

345

REFERENCES

1. E.H. Dowell, AlA A J., 8(1970), 385-399. 2. M.D. Olson, AIAA J., 8(1970), 747-752. 3. J.N. Rossettos and P. Tong, J. Applied Mech., ASME, 41(1974), 1075-1080. 4. J.W. Sawyer, J. Aircraft, 14(1977), 387-393. 5. R.W. Clough and C.A. Felippa, Proc. 2nd Conf., AFB, 1968 .

• Plx.y.t)

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OPTIMISING THE GEOMETRY OF ENERGY ABSORBING COMPOSITE TUBES WITH PARTICULAR REFERENCE

TO RAIL VEHICLE APPLICATION

J.F. KELLY

British Rai/ways Board Research Division - Railway Technical Centre

London Road - DE2 8UP DERBY - England

Rail vehicle structural bodyshell designs are usually welded fabrications of aluminium or steel, the plasticity of which dominates the energy absorption when structural collapse occurs during severe frontal impact. Composite materials have been shown to offer significant increases in energy absorption, when compared with metals during controlled collapse regimes. This experimental work confirms the advantages of such materials, at force and energy levels compatible with rail vehicle design constraints and identifies the influence of geometry on cylindrical energy absorbing modules manufactured from glass fibres and polyester resin.

1. INTRODUCTION

Metals, used for the fabrication of rail vehicles - aluminium or mild steel -generally possess considerable plasticity and it is this property which provides the energy absorption in such structures during impact. If the structural collapse is allowed to initiate and progress in an unrestrained manner, the structure can fail catastrophically with low energy absorption. If the collapse is initiated at a particular point, that is ''triggered'', and allowed to propagate in a stable manner, the level of energy absorbed can be substantially raised (1).

Within the spectrum of practical structural metallic shapes, the most advantageous method of absorbing energy during structural collapse is by bending, and the most effective shape of a metallic component is a cylindrical tube, arranged to collapse in an axial mode. The load/deflection curves for such collapsing tubes show a peak load at which plastic buckling commences, followed by an oscillating load cycle generally at a much lower level than that of the initial peak load.

348

The materials generally considered for thermosetting fibre reinforced plastic (FRP) composite structures comprise fibres of glass. graphite or Kevlar and a resin matrix usually of epoxy. phenolic or polyester. FRP structures therefore. possibly with the exception of Kevlar. are considered to be of a brittle nature. However. it has been reported that small cylindrical tubes manufactured from combinations of these materials can be triggered into a stable collapse regime which tends to occur at a constant load. without the large initiation loads seen with metal tubes (2). The characteristic of the collapse results in a much enhanced level of energy absorption. When compared with the metallic tubes. this enhancement is in part brought about by the creation of fracture surfaces. rather than by plastic deformation and in addition. a much larger proportion of the tube length can be used more effectively.

2. PROBLEM DEFINITION

When considering absorbing the kinetic energy of a rail passenger vehicle. or a complete rake of vehicles. during frontal impact. the first requirement would be to limit the deceleration to a value which the passenger could tolerate without unacceptable injury. The second requirement would be to limit the force level in order to protect the main structure of the vehicle. The level of acceptable deceleration. dictated by passenger safety. would then determine the distance over which the force could be applied. and ultimately the amount of energy which could be absorbed.

The mass of a typical inter-city rail passenger train. arranged as a fixed rake of vehicles would be some 600 tonnes and at a maximum service speed of 200 kmlhr represents a kinetic energy of 0.93 GJ. It soon becomes apparent that. based on the criteria outlined above. one could only sensibly attempt to absorb energies at the lower end of the speed range and our studies have indicated that this may be economic. based on the savings in repair costs alone (3). Part of the effort therefore has been directed towards protecting the critical parts of the forward facing structural components. together with areas such as vehicle interconnection.

The maximum compressive force level to be resisted by a British Rail passenger vehicle. without permanent deformation. is 2 MN. Therefore. any device would be required to 'trigger' and maintain stabilised collapsing forces below this value. Previous damage analysis of passenger vehicles had indicated that a sensible velocity range would be 4 to 14 m/s.

349

3. TEST SPECIMENS

In order to maintain reasonable quality control of the dimensions of the samples and the resin/glass ratio during manufacture it was decided to utilise the filament winding technique. The materials of construction were dictated by the economic constraints of the application and against this background continuous filament "En glass fibres were selected together with a polyester resin of the type Crystic 491. The range of geometric dimensions of the test specimens was selected against the likely installation envelope available on the vehicle and an approximation of achievable energy densities previously reported (2). Previous workers (4) had suggested that a chamfer on the end of a composite tube was sufficient to trigger a stable collapse regime. It was therefore decided to machine a 45° chamfer on one end of all the test specimens, Fig 1. The target glass/resin ratio for all the samples was 75%.

4. TEST VARIABLES

From the foregoing, the test variables selected were:-

Impact velocity - Quasi-static Impact velocity - Dynamic Impact energy Tube inner diameter Tube wall thickness Tube length/diameter ratio

5. TEST EQUIPMENT

1 mm/sec 4 to 14 m/s up to 300 kJ 5 to 300 mm 10 to 20 mm 1:1and2:1

The testing was carried out to compare the quasi-static and dynamic modes over the variables listed above.

5.1 Quasi-Static

The quasi-static testing was conducted on a DARTEC 2000 kN Universal Testing Machine, Fig 2. The force/deflection relationship of the test specimen was recorded on a two axis plotter and the test velocity, which is the velocity of the closing plattens, was kept constant at 1 mmlsec throughout the tests. The machine had been previously calibrated to British Standard 1610, Class 1, indicating an error of less than 1%.

5.2 Dynamic

The dynamic testing was conducted on the British Rail Dropped Weight Test Facility Fig 3, which has a maximum capacity of 300 kJ and a maximum drop height of 10m. The load cell, was calibrated at the National Engineering Laboratory and the general accuracy of the installed instrumentation is better

350

than +/- 2% with the exception of the values of acceleration, which were +/-7%. Timing slots were attached to the tup which, with the aid of a laser, enabled the position and velocity to be determined within an interval of 5 mm.

The dynamic tests were conducted using a mass of 3 tonnes; the drop height was varied from 0.8 to 9.0 m resulting in an impact approach velocity within the range 3.8 m/sec to 13.3 m/sec. Due to the presence of high frequency components in the force recordings the effect of filtering was investigated and it was decided that filtering at 1024 Hz was most acceptable. The main effect of filtering is to reduce the force peaks with very little effect on the energy levels.

6. TEST RESULTS

6.1 Quasi-Static Compression

The maximum force recorded during the collapse of the tubes ranged from 252 kN to 1650 kN whilst the mean force - the value used for calculating the absorbed energy - ranged from 184 kN to 1240 kN for the 75 x 75 x 1 0 mm and .300 x 300 x 20 mm tubes respectively. Repeatability of the "triggering" force was investigated over five samples and found to be +2.5% to -3% about the arithmetic mean for the 300 x 300 x 20 mm tubes and +8.5% to -10.5% for the 150 x 150 x 10 mm tubes.

The 300 x 300 x 10 mm tubes exhibited a sharp force peak as the collapse was initiated, resulting in a peak/mean ratio of approximately 1.9 whilst as the wall thickness was increased to 20 mm the ratio fell to 1.33, with a much less sharp characteristic. As the tube diameter was reduced to 150 mm the force peak was much less defined at the 10 mm or 20 mm wall thickness and the peak/mean force ratio fell to approximately 1.1. A reduction still further to 75 mm diameter produced an even smoother transition from the onset of collapse to the stablized condition both at the 10 mm and 20 mm wall thickness.

The energy density was calculated from the mean force extracted from the force/distance trace and the measured length of crush, adjusted in volume for the initial taper. The highest energy density obtained was 52.8 kJ/kg using the 75 mm x 75 mm x 20 mm tube and the lowest at 23.5 kJ/kg obtained from the 300 x 300 x 10 mm tube. As the LID ratio of the tubes was increased to 2 a significant reduction in energy density occurred - of the order of 25% - which was characterised by a very sharp drop in the force trace and shear bands appearing on the tube walls remote from the crush zone.

6.2 Dynamic Compression

In view of the relatively poor performance of the high UD ratio samples during the quasi-static tests it was decided to delete this configuration from the dynamic tests. The maximum force occurring during collapse ranged from

351

339 kN to 2278 kN for the 75 x 75 x 10 mm and 300 and 300 x 20 mm tubes respectively, each at an impact velocity of 8.3 m/sec. The range of peak/mean force ratio stayed sensibly constant compared with the quasi­static results except for the 300 x 300 x 20 mm tube which peaked at 2.29 at 8.3 m/sec. The energy densities, which were calculated using the mean force from the force/distance traces filtered at 1024 hZ over the distance crushed and correcting for the taper, ranged from 22.7 kJ/kg to 58.2 kJ/kg for the 300 x 300 x 10 mm tube at 13.3 m/sec and the 75 x 75 x 20 mm tube at 3.8 m/sec respectively. Fig 4 shows a typical force/disp. trace filtered and unfiltered.

7. DISCUSSION

The 45 deg taper is very successful at initiating the collapse at these high force levels over the range of tube geometries tested in both a quasi-static and dynamic mode. The indications are that as the UD of the tube is increased the collapse will still initiate at the tapered face although shear failures will occur remote from this which considerably reduce the absorbed energy. The effect of wall thickness/diameter ratio on energy density is unclear. From the quasi-static tests the energy density remains sensibly constant against TUD2 down to a value of .05 although such a correlation was not apparent from the dynamic tests. From Fig 5 it can be seen that the energy density has a tendency to reduce above some 4 m/sec and this reduction tends to be greater with the smaller diameter tube. At the 300 mm diameter tube the energy density was virtually unaffected by velocity.

8. CONCLUSIONS

CompOSite materials, in the form of glass/polyester filament wound tubes, can offer significant advantages for energy absorption in rail vehicle end structures and interconnections. They have a high energy density relative to most metals and they can be manufactured cheaply. The degree of this advantage is indicated in Fig 6 where the composite tube illustrated has an energy capacity at least equal to the energy that was used in statically crushing the vehicle end.

The Author would like to thank the Directors of British Railways Board Research Division for permission to publish this work and colleagues within the Division who gave invaluable assistance, together with Prof. D. Hull of the University of Cambridge for his helpful comments.

REFERENCES

1. Coppa, A.P., NASA TN-D1510, (1962) 2. Magee, C.L., Thornton, P.H., S.A.E. 780434 (1978) 3. Scholes, A., I.Mech.E. Conf. Railway Vehicle Body Structures., C284/85 :

Derby 1985 4. Kirsch, P.A., Jahnle, H.A., Soc. Auto. Eng. (1981) No 0148-7191/81/0223-

0233502.5.

352

ElQ.1 SAMPLES BEFORE/AFTER IMPACT fKU. QUASI·STA TIC TEST MACHINE

DISPLACEMENT

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~ VARIATION OF ENERGY PENSITY WITH VELOCITY

EIQ.2 COMPARITIVE ENERGY

ABSTRACT

STRESSES IN THE JOINT OF AN END FITTING TO A COMPOSITE TORQUE TUBE

H. BROWN, R. HAINES', T. JOHNS, J. MURPHY"

Portsmouth Polytechnic - Mechanical Engineering department 'Portsmouth Polytechnic - Mathematics Department Anglesa Road - POI 3DJ PORTSMOUTH - England

"High Temperature Engineers Ltd Lower Ouay Road - Fareham - P016 ORO PORTSMOUTH - England

When an end fitt i ng is bonded to a composite tube it is normal practice to use thick joint theory to estimate joint shear stresses. Because a composite is much stiffer in the longitudinal than the through thickness direction, the whole tube acts as the joint. Thus the peak stresses depend on the geometry of the joint and the material properties. Mathematical analysis and finite element analysis illus­trate stress curves. Consideration is given to an end fitting comp­rising a mandrel and cuff-nut with mechanical locking.

1. TORQUE TRANSFER BETWEEN TUBES

The transfer of load from one structural member to another is often complex. For analysis, consideration must be given to equilib­rium of forces and to compatability of displacements. In the case of bonded joints, a critical consideration is the thickness of the joint. Whilst glues generally perform best in very thin layers, stress analysis shows that the stresses in such a layer are not uniform and are usually extremely high at the ends of the joint. For fibre reinforced composites, a thick joint approach may be used because of the presence of the resin in the composite.

Consider two tubes joined together such that one fits inside the other. A torque is app 1 i ed to the end of one tube, transferred through the joi nt and reacted at the end of the second tube, The most elementary theory that may be used is compound tube theory; the torque at any cross-section is shared by the components in the ratio of tor­sional stiffness, GJ. At changes of cross-section, transfer of torque is a step function, requiring infinite shear stresses at the joint.

354

When the tubes are isotropic materials, the model is accurate except very close to the junction. Practical components should include methods of reducing the high shear stresses.

A composite material is orthotropic. In a tube designed for torque transmi ss i on, fi bres are predomi nant in the ± 45 0 di recti on. The material is thus much stiffer in the plane of the laminate than through the thi ckness. Indeed, the whole thi ckness of the materi a 1 may participate in the torque transfer process.

2. MATHEMATICAL ANALYSIS

Consider two tubes joined together with one inside the other as shown in Figure 1. Coordinates r radial, e angular and z axial locate a point within a tube. An applied torque is transmitted axially by the shear stress distribution Tze, the distribution of which is changed by the presence of the shear stress Tre' The stresses on an element, Figure 1, provide equilibrium forces and the equilibrium equation is:

T re

aT + r_re

ar = 0 ..... (1)

If , is the angular rotation of a point, Figure 1, the shear strains must be compatible, i.e:

a, y = r - and y

re ar ze

a, r­

az ..... (2)

Equations (1) and (2) are connected by the material shear compliances

= STand y = S T re re ze ze ze

Combining (1) (2) and (3) yields the equation

a2, 2 a, a2, + - - + k2 .- = 0

ar2 r ar az2

where, 2

k Sre S

ze

A solution to (4) is:

, = (Asinxr + Bcosxr)(ce~ + D;~) + Ez + F

..... (3)

..... (4)

..... (5)

355

the boundary conditions determining the constants.

Figure 2 illustrates the most simple problem of a tube being joined to a rigid support. The boundary conditions are shown in Figure 2. The external boundary conditions dictate a series of >..

va 1 ues dependant on the rat i 0 of di ameters. The constants for each >.. value must be such that. the cOlll1lon boundary conditions at the junction are satisfied. The redistribution of TZ e from the linear cOlll1lences away from the joi nt, and the peal< value of joi nt shear stress depends on the geometry of the joint and the value of 1<.

More complex problems necessitate numerical methods of solu~ion. The PAFEC software i ncl udes model s for lami nates and orthotropic materials within the finite element solution.

3. DESIGN FEATURES AND DATA

The basic geometry of a joint may be determined by the compound tube theory and will consi st of the di ameters of components. An average joint shear stress may be calculated on the basis of the simple joint shown in Figure 3. The curves are the variation of joint shear stress along the length for a I< value of 2. They clearly show that the peal< joint stress is affected by the tube thicl<ness.

Peal< joint shear stresses may be reduced by decreasing the st i ffness of the inner tube at the end. Thi s may be achi eved by an angle, the effect of which are illustrated in Figure 4. This is simply the application of good design practice of avoiding sharp internal corners.

In the case of an end fitting, a taper may be used to gradually increase the torsional stiffness and distribute the load transfer over a longer length of the tube, thus reducing the peal< joint shear stress.

A practical end fitting, Figure 5, may comprise a mandrel, fitting inside the tube and a cuff-nut, fitting outside the tube. Such an arrangement provides a neat compact fitting. A typical arrangement mi ght have the mandrel extendi ng further into the tube than the cuff -nut extends over it. For such an arrangement, the torque is initially transferred partially from the tube to the tube and mandrel and then proportioned to the tube, mandrel and cuff -nut. Finally it is transferred to the mandrel and cuff-nut which are rigidly connected. A finite element analysis quantifies the peak stresses and distribution of stresses.

4. MECHANICAL LOCKING

Because of the large joint shear stress at the end of a bonded joint, many designs incorporate mechanical locl<ing features to in­crease confidence in the design. This feature is attractive, since it provides a fail-safe design and improves the joint by providing more uniform stresses in the jOint.

356

Mechanical systems passing through the laminate, such as rivets and bolts, require careful consideration about reinforcement, etc.

A popular mechanical locking system is a shaped mandrel. The shape of the mandrel increases the joint shear stress at a cross­sect i on but thi sis more than compensated for by the more uniform joint shear stress distribution. The increase in joint shear stress due to shaping is illustrated in Figure 6, a factor of 1.5 being indicated.

For the mechanism to be fail safe, the mechanical locking mech­anism should be sufficient in itself to transfer the loading. It must be assumed that the interface loses both its shear and tensile stress carrying capabilities. Relying solely on compressive stresses at the interface, the tube and mandrel are subjected to bending round their circumferences. Should their bending stiffnesses be insuf­ficient the tube may deflect and rotate about the mandrel as shown in Figure 7. A cuff-nut provides additional protection against this defl ect ion.

5. FURTHER CONSIDERATION

The practical considerations of construction of the joint fitt­ing are of technical and commercial importance. To maintain an ad­vantage over their metal counterparts, weight and cost must be kept to a minimum, whilst confidence in performance must be established.

Winding Techniques are such that it is now possible to integ­rally wind flanges on a tube. Such developments should improve the attractions of the composite tube. 5'

/

/ /

f-

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FIG. 1. PARAMETERS, STRESSES AND STRAINS

357

<P'~OCOMMON ------

FIG. 2. BOUNDARY CONDITIONS FOR TUBE TO RIGID CONNECTION

6-

5

4

T ~ .----~-.. _.25_D ---, .. ,

g T ~3 <;)

),< L __

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x/D FIG. 3. EFFECT OF TUBE THICKNESS ON Tre

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o o .05 .10 .15 .20 x/D

FIG. 4. EFFECT OF TAPER OF MANDREL ON Tre

358

V) V) w c:: l­V)

1

TUBE

FIG. 5. COMPONENTS OF FITTING

- POSITION ON ·MANDREL

FIG. 6. JOINT SHEAR STRESS

FIG. 7. MECHANICAL LOCKING FORCES

-- NOMINAL SHEAR STRESS

STRENGTH AND RESPONSE OF COMPOSITE PLATES CONTAINING AN OPEN HOLE AND SUBJECTED

TO COMPRESSIVE LOADING

ABSTRACT

F.-K. CHANG, L. LESSARD

Stanford University - Department of Aeronautics/Astronautics 94305 STANFORD - USA

A computer code, designated as "PDHOLEC", was developed for analyzing compression failure of laminated composites containing an open hole. The code was developed based on the progressive damage analysis which was proposed by the authors. The analysis consists of a stress analysis for calculating stresses and strains inside the laminates, and a failure analysis for predicting damage and determining the residual stiffnesses and strengths of the laminates. Not only can the ultimate load of the laminates be determined by the code, but also the types and the extent of damage as a function of the applied load can be evaluated. The code was verified extensively by experimental data. The code can be used as a design tool for designing and sizing laminated composites containing holes and subjected to compression loading.

I. IJ'Io'TRODUcnON

Tensile strength of laminated composites containing holes or cutouts have been studied extensively in the literature. Both analytical and experimental investigations have been performed and several analytical models have been developed for predicting the tensile strength of laminated composites containing holes and cutouts. However, there is very limited information available in the literature about compressive strength and re­sponse of laminated composites containing holes and cutouts. Most of the published work has been experimental.

Since more and more composite structures are designed for primary load-bearing elements, the knowledge of the ability of these structures to sustain compression becomes increasingly important and is essential for the safety design of the structures.

Recently, a progressive damage model was developed for evaluating compression failure of laminated composites containing holes. A computer code. designated as

360

"PDHOLEC", was developed based on the model (see Figure 1). The code can provide the following information:

1. the types and the size of in-plane damage at a given load. 2. the residual stiffness and strength of the laminates. 3 . the final collapse load. 4. the complete response of the laminates from initial loading to final

catastrophic failure.

The objective of this investigation is to perform a parametric study to evaluate the effect of geometry, ply orientation, and loading direction on the compression response of laminated composites containing holes.

n. COMPU1ER CODE

The "PDHOLEC" code, which is based on a progressive damage model, was developed by the authors /1/. Basically, the model consists 9f a stress analysis and a failure analysis. Stresses and strains in the plates were calculated by a nonlinear finite element analysis, which is based on a finite deformation theory with consideration of both material and geometric nonlinearities. Plane stress condition was used for the development of the analysis.

The types and size of damage were predicted by a proposed failure analysis which includes a set of proposed failure criteria and property degradation models for each mode of failure. In-plane failure modes such as matrix cracking, matrix compression, fiber­matrix shearing, and fiber buckling (kinking) were the basic failure modes considered in this investigation. The effect of stacking sequence on the strength and failure was taken into account in the analysis, although ply delamination and the associated ply buckling as a result of stacking sequence effect were not considered.

Basically, stresses and strains are calculated at each incremental step, and evaluated by the failure criteria to determine the occurrence of failure and the mode of failure. Mechanical properties in the damaged area are reduced appropriately according to the property degradation models. Stresses and strains will then be recalculated to determine any additional damage as a result of stress redistributions at the same load. This procedure will continue until no additional damage is found, and the next increment is then pursued. The final collapse load is determined when the plates cannot sustain any additional load. Detailed derivations of the model are given elsewhere /1/.

ID. COMPARISONS

Extensive comparisons have been made between the numerical simulations obtained from the code and the data generated from the tests. Seven different ply orientations and more than forty various geometries were considered. Details of the comparisons and the test results are given in /1,2/, hence only some typical results of comparison are presented in the following.

361

Figure 2 illustrates the predicted types and size of damage inside a [+30/-30]5 specimen as a function of the applied load. Damage was ~redicted near th~ stress concentration areas and was confined to local areas up to the pomt where the apphed load was approaching the final collapse load. The damage then grew suddenly in un­controllable fashion and propagated along 30 degree lines, measured from the loading direction to the free edges. The specimen failed catastrophically. It is worth noting that the failure mode was dominated by matrix tensile cracking. The predicted damage pat­tern was very similar to the test results, as shown in Figure 3.

For [(0/90)6]8 laminates, the results of the model's prediction are also presented in Figure 4 at different load levels. Matrix cracking initiated from the hole boundary near the stress concentration area and then propagated along the loading direction (parallel to the 0 degree fibers). The specimen continued to sustain additional load, until fiber buckling (fiber kinking) was predicted in the 0 degree plies near the hole boundary. Fiber buckling grew quickly in the direction normal to the loading direction and resulted in the total collapse of the specimen. The predicted damage pattern also coincided with the test results /1,2/.

The effect of ply clustering (plies with the same ply orientation grouped together) on the strength and damage of the notched laminates can also be studied by the code. For example, considering [06/'906]5 laminates, the predicted damage pattern is shown in Figure 5. Unlike [(0/90)6]5 laminates, damage, dominated by fiber-matrix shearing, originated from the hole boundary and grew in the direction parallel to the loading direction. The predicted damage size was much more extensive than that predicted in [(0190)6]8 laminates. Such extensive damage would most likely induce delaminations, because delaminations could be initiated by in-plane failure such as matrix cracking, fiber matrix shearing, and fiber buckling.

N. CONCLUDING REMARKS

A computer code, "PDHOLEC", was developed for predicting in-plane response and compression failure of laminated composites containing an open hole. The code can be used as a design tool for designing and sizing laminated composites containing holes subjected to compression. The information provided by the code include compressive strength, in-plane response, types and extent of in-plane damage in each layer in the laminates. However, the code does not consider Euler buckling of the plates and ply buckling associated with delamination.

Due to limited space, only some of the results are presented in this paper. More information can be easily generated by use of the code. The code can be obtained from Fu-Kuo Chang at the address given in the front page.

~. ACKNOWLEDGEMNET

The support of the Charles Lee Powell foundation grant and the F.C.A.R. of Quebec, Canada fellowship for this investigation are gratefully appreciated.

362

VII. REFERENCES

1. Chang, F. K. and Lessard, L., Compression Failure of Laminated Composites Containing an Open Hole-Analytical Method, J. of Composite Materials, (submitted).

2. Chang, F. K. and Lessard, L., Compression Failure of Laminated Composites Containing an Open Hole-Experiment, 1. of Composite Materials, (submitted).

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INTEGRALLY WOVEN SANDWICH·STRUCTURES

ABSTRACT

K. DRECHSLER, J. BRANDT", F.J. ARENDTS

University of Stuttgart Pfaffenwaldring 31 - 7000 STUTTGART 80 - West Germany

"MBB GmbH - 8 MUNCHEN 80 - West Germany

New technical opportunities with regard to manufacturing and me­

chanical properties are opened up by the development of new int.,­

grally woven sandwich structures. In these fibre constructions, two

fabric layers are separated by a woven-in system of z-directional

fibres. After impregnation with a resin, the double-wall fabrics form

the sandwich structure. The mechanical properties can be adjusted

in a wide range through the number and the arrangement of the

thread-linking system. A comparison with foam and honeycomb sand­wich structures shows the potential of this new material class.

INTRODUCTION

The development and application of new textile technologies, such

as 3-D weaving, 3-D braiding or knitting have shown that there is a

great potential for decreasing manufacturing costs and to impro­

ving mechanical properties of composites [1,2,3l.

First of all, the 3-dimensional reinforcement has certain advantages

with regard to interlaminar shear strength and fracture behavior.

It could be shown that 3-~ composites are more resistant to damage

because the fibres in the 3rd direction will absorb some of the impact

energy. An application is the use of new textile technologies for

manufacturing core materials for sandwich structures. For example,

366

the use of braided tubes or deep-drawn knitted fabrics leads to interesting properties [4]. Figure 1 schematically shows some of

these structures.

It is characteristic for all of the sandwich structures shown that the skins are manufactured separately and bonded to the core in a second step. This, in addition to the relatively costly manufacturing, leads to disadvantages in some mechanical properties. Due to this construction, there is a tendency to a debonding of the skins from the core during bending, buckling or impact loads. To guarantee a better connection between the top and bottom skins for this appli­cation, a 3-D fibre structure is also desirable. It is shown in the following text that manufacturing these fibre structures is possible by developing a new textile process that allows totally integrated sandwich units to be woven in one step.

1. Manufacturing 3-D Sandwich Structures

In Figure 2, some types of the 3-D fibre structures developed are shown schematically. The length of fibres in thickness direction, the height of the sandwich, as well as the arrangement of the linking threads can all be varied in a wide range. Through this, both stiff­ness and strength properties as well as the core density can be adjusted to the respective requirements. The present developments have been focused on the application of glas fibres, but aramide and carbon fibres can al so be used.

Making up double-wall fabric to obtain a consolidated sandwich is very easy. The fabric is impregnated, for example, with an epoxi resin and the excess resin is squeezed out. The linking fibres between the walls have such a high stiffness, caused by the textile process, that they stand up after the squeezing process. This leads to a hollow structure with two impregnated fabric layers and impregnated linking fibres after curing. Because no pressure is used during curing, only the bottom fabric bec~mes smooth. On the top wall, the surface of the impregnated woven structure remains.

With this textile technique, only flat or slightly curved sandwich panels can be manufactured. If the desired radius of curvature is too small, no constant thickness can be realized. For this special application, knitted double-wall structures were developed. In this case, the walls consist of knitted fabrics, which are also connected by worked-in fibres.

367

2. Testing Program

To evaluate the mechanical properties, the following experiments

were carried out:

- compression tests

- flatwise tensile tests

- four-point bending tests

- s hear tests

(DIN 53291)

(DIN 53292)

(DIN 53293)

(DIN 53294)

Besides this, the buckling and impact behaviour was examined quali­

tatively.

For comparison, respective experiments were carried out using con­

ventional sandwich structures with honeycomb (NOMEX ECA 3.2-48)

and foam (ROHAZELL 51A) cores. The core density of these materials

amounted to approximately 50 kg/m 3 , the core height was chosen

according to the respective integrally woven structures (Smm,

7.5 mm,10mml. For the skins, glas fibre satin weavings (163g/m 2)

bonded to the core were used. The 3-~ sandwich structures tested

also consisted of glass fibres.

In Figure 3 to Figure 6 some results of the mechanical tests are

summarized.

The test results show the performance of the different sandwich

structures depending on the core density. The values of the in­

tegra"y woven materials represent the state of the development up

to now. Further improvements can be expected by optimising the

manufacturing process and the respective choice of fibres and

matrix materials.

Figure 4 and Figure 5 show that the range of the density the thread­

linking system realized up to now amounts to 50 to 100 kg/m3 . The

height of the 3-D sandwich structure can be varied from 5 to 10 mm.

With both figures, the large range of obtainable mechanical values,

depending on the fibre structure, is illustrated. A compression

strength ranging from 0.3 to 3.1 N/mm2 and a shear strength from

0.3 to 1.6 N/mm2 is achievable. These values demonstrate the good

performance of these structures in comparison to comparable (with

repect to density) honeycomb and foam sandwich structures.

The shear stiffness of the 3-~ structure is mark able lower than

368

those of honeycomb but comparable with those of foam structures.

The best material with respect to flatwise tensile strength is the

integrally woven sandwich structure.

The direct comparison of bending, impact and buckling perfor­

mance is difficult due to the high influence of the used skins. The

kind of failure behaviour is of more interest (see Figure 71.

The most significant property of the integrally woven material is the

good damage tolerance behaviour being expressed by the fact that

the failure is locally limited. Especially no delamination of the skins

is possible.

3. Conclusions

By developing and applying a new textile process to advanced fibres,

it was possible to manufacture integrally woven sandwich struc­

tures with interesting properties.

Furthermore, knitted 3-D fibre structures have been developed

allowing spherical curved sandwich panels to be manufactured. For

many applications, a special interesting feature is the fact, that the space between the skins is not totally filled. The more or less hollow

structure allows this space to be used, for example, for cable laying.

The mechanical properties can be adjusted in a wide range for different

applications. They exceed partially those of the foam and, in tension,

even those of honeycomb sandwich structures. An interesting feature

is the good damage tolerance behaviour of these structures.

By further optimising of manufacturing, fibre arrangement and the use

of carbon fibres, additional improvements of mechanical properties

can be expected.

ACKNOWLEDGEMENT

This work is the result of a cooperation between the Central Labora­

tories of M88 GmbH, Munich, Vorwerk& Co, Kulmbach and the Institute

for Aircraft-Design, University of Stuttgart, and was supported by

the German Ministry of Research and Technology

369

REFERENCES

1 - Drechsler K., Der EinfluB der Faserstruktur auf die Steifigkeit

verstarkter Kunststoffe, DGLR Symposium, Berlin (1987)

2 - Arendts F.J., Drechsler K., Verbesserung der Eigenschaften fa­

serverstarkter Kunststoffe durch den Einsatz neuer Textiltech­

nologien, Verbundwerk 88, Wiesbaden (1988)

3 - Bottger, Drechsler, Siegling, Anwendung von dreidimensionalen

Faserstrukturen bei Verbundwerkstoffen mit Polymermatrix,

BMFT- Symposium Materialforschung, Hamm (1988)

4 - Williams, High-Performance Sandwich Panels made from Braided

Tubes, 33rd International SAMPE Symposium, Anaheim (1988)

5 - Drechsler K., Wierse M., Experimentelle Untersuchung von Sand­

wichstrukturen mit verschiedenen Kernmaterialien, Internal

Rep 0 r t, Ins tit ute for Air cr a ft Des i 9 n (1 9 88 )

CDre Mat"ia/, 'Dr Sandwich Structure,

Foam HoneJcomb

Braided Tube. Deep-Drawn Knitted Fabric

Figure

370

Int'grall, W'rln Sandwich Stuetu",

Type 1 Type 2 Type 3

Figure 2

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ELABORATION PROCESSING

Chairman: Dr KELLERER M.B.B.

THE DEVELOPMENT AND APPLICATION OF THE MULTI LIVE-FEED MOULDING PROCESS FOR THE PRODUCTION

OF INJECTION MOULDINGS CONTAINING LAMINATED AND OTHER SPECIFIC FIBRE

ORIENTATION DISTRIBUTIONS

P. ALLAN, M.J. BEVIS

Wolfson Centre - Brunei University Kingston Lane - UB83PH UXBRIDGE MIDDLESEX - England

ABSTRACT

The extension of the Multi Live-feed Moulding (MLFM) process to four independently controlled live feeds is described. Three mould cavities, a rectangular plaque, a tapered fin and a circular ring are used to illustrate the degree of control over the orientation of fibres that is possible using the MLFM process. Contact microradiography and mech3nical test results are used to illustrate the enhancement of fibre orientation control and mechanical properties that is possible with injection moulded glass fibre reinforced thermoplastics.

INTRODUCTION

Multi live-feed moulding (MLFM) (1,2) provides a route for the management of fibres in complex shaped moulded parts. The process is based on the use of independently controlled melt pressure sources at a multiplicity of gates to a mould cavity for the purpose of controlling microstructure in short fibre reinforced plastics. The process as originally described (1) was based on the two live-feeds. This demonstrated substantial enhancement of Young's modulus and internal weld line strength that could be gained from the use of the MLFM process.

The MLFM process has been developed further to enable the production of injection mOUldings containing laminated and other specific fibre orientation distributions.

376

FOUR LIVE-FEED MOULDING

One embodiment of the four independently controlled live­feed arrangement utilises an injection moulder fitted with twin conventional injection units, each providing melt for a double live-feed device. The schematic diagram (Fig.l) illustrates the use of two double live-feed devices (D). On completion of initial ~ould filling of the mould cavity (C), and in place of the conventional packing phase, the pistons 1-4 are caused to oscillate anG thereby influence fibre orientation during solidification of the melt. The influence of the operation of live feeds on fibre orientation is illustrated below. The operation for the independent control of the pistons 1-4 is governed by a microprocessor, which is additional to the control system for the twin injection machine.

The operational four live-feed machine in the Wolfson Centre for Materials Processing at BruneI University is based on a Negri Bossi 130-90-90 two colour injection moulding machine. The two double live-feed devices were designed, constructed and integrated with the NB 130-90-90 at BruneI University. The microp~ocessorcontrol was built to specification by Technology Concepts.

2 APPLICATIONS OF THE MLFM PROCESS

Three applications of the process for the moulding of (i) square plaques, (ii) tapered fin and (iii) circular ring ~ilil

square cross-section, ~re used here to illustrate the mode of operation. All three mould cavities were fed from the same four runner geometry, terminating in the four gates (1-4) illustrated in Figure 2a. Characteristic features of the three cavities, the orientation of reinforcing fibres and the mechanical properties of the moulding are reported below.

2.1 Square plaques

Figure 2a is a schematic diagram of a square plaque cavity with four independently controlled fan gates. On completion of initial mould cavity filling, the pistons servicing gates 3 and 4 are caused to move back and forth and 180 0 out of phase. This results in a preferential alignment of reinforcing fibres at the melt-solid interface and parallel to 3-4 in Figure 2b. After a pre-set time the pistons servicing gates 1 and 2 are caused to operate out of phase and in place of pistons 3 and 4, thereby causing a layer of fibres with a preferred orientation parallel to 1-2 in Figure 2c to be formed. A multi layer laminated structure may be formed through the thickness of the plaque

according to the sequencing of the operation of any pair of live feeds as the material solidifies from the outside to the centre of the plaque cavity. Figure 2d indicates the position from which contact IT.icroradiographs were taken.

Figure 3 is a contact radiograph showing a laminated structure 30% glass reinforced polypropylene produced in a 6mm thick moulding. The radiograph represents the changes in preferred orientation from the surface ( S) to the central core (C) of the moulded plaque. A laminated structure is apparent where the preferred orientation of the fibres in the plane of the radiograph is represented by (--) and normal to the plane of the radiograph by O. The alignment of the fibres over an extended area in the core was caused by deliberate extension in time of the out of phase operation of pistons 1 and 2.

2.2 Circular ring

The l50mm internal diameter ring with a 10x15mm cross­section was gated as illustrated in Figure 4a. Initial mould filling was with 30% glass fibre reinforced polypropylene through a single gate (gate 1) for all mouldings referred to below. This produced an initial internal weld line opposite gate 4. Mechanical test data is given in Table 1 for mouldings produced with:

1) a static packing pressure

2) Oscillating packing pressures applied through all four gates

3) Multi live-feed packing with the sequencing of piston movements at the four gates as illustrated in Figures 4b and 4c. Alternate sequencing in accord with Figures 4b and 4c eliminated the internal weld line and provided for the optimum circumferential alignment of glass fibres.

Table 1 shows that the use of MLFM results in an increase in stiffness of :: 28 per cent, and an increase in weld line strength of :: 70 per cent.

2.3 Tapered fin

The geometry of the tapered fin ranging in thickness from 2mm to 5mm is illustrated in Figure Sa. In the example selected for presentation the initial injection of 30% glass fibre reinforced polyarylamide (IXEF) was through gate 4 as illustrated in Figure 5b. The MLFM sequence illustrated in Fig. 5d was used to impart enhanced transverse orientation of glass fibres. The contact X-ray radiographs shown in Figures 5c and 5e illustrate

377

378

the difference between the preferred orientation of glass fibres through the thickness of conventional and MLF mouldings. Transverse orientation of fibres to the outer surface of the moulding is evident in Fig. 5e. Three point bend tests carried out parallel to gates 2-3 showed, in the example illustrated,that the application of MLFM resulted in a 22% increase in flexural stiffness. This corresponded with 19% decrease in flexural stiffness for tests carried out parallel to gates 1-4.

3 CONCLUDING REMARKS

The three examples of mould geometries employing multi live­feed presented above, illustrate a control over preferred orientation of reinforcing fibres and mechanical properties not usually possible with conventional injection moulding. These should serve t indicate the general application of MLFM for the management of t~bres in complex shaped mouldings. More detailed presentations based on the application of MLFM for a range of mould cavity geometries and fibre reinforced thermplastics, thermosets and liquid crystal polymers are being prepared for publication (3).

ACKNOWLEDGMENTS

The authors are indebted to the British Technology Group and British Aerospace for financial support.

REFERENCES

1. P S Allan and M J Bevis. Plab.lcs and Rubber, Processing and Applications, 7 (1987) 3-10.

2. P S Allan and M J Bevis, UK Patent Application 85-31374.

3. P S Allan, M J Bevis and co-workers (in preparation).

Table 1 Relative tensile modulus and failure load measured at positions A, Band C (Fig. 4B).

Position Static OPP MLFM (single gate) single gate fill single gate fill

4 gate ~ack

Relative I Failure Relative Failure Relative Failure modulus load(KN) modulus load(KN) modulus load(KN)

A 1.04 5.45 1.02 6.3 1.29 7.33

B 1.03 4.13 1.00 6.8 1.29 7.6

C 1.02 3.6 1.01 5.7 1.29 6.1

Fig.l - Schematic diagram of a four live-feed arrangement

o

o

o

Fig.3 - Contac t radiograph showing laminated microstructure

Fig. 4a ~ The circular ring moulding

3 2a

~ 3

2b

3

2c

2d

379

4

, 2

c ~-o

Fig.2 - Square plaque mould with four independently controlled gates. Diagrams 2b and 2C show their mode of operation. The microradiograph in Fig.3 corresponds to the section illustrated in Fig. 2d.

380

Figs. 4b and 4c-Showing the MLFM operating sequences.

Sb

i t t 4

Sc!

4

Sc

Sa Fig. Sa -Tapered fin · moulding Figs. Sb/ c - Conventional injection moulding through gate 4 and corresponding radiograph. Section parallel to 1-4 and normal to diagram.

Figs.Sd/e -MLFM through gates 2 and 3.

FABRICATION OF FIBER REINFORCED CERAMIC COMPOSITES

ABSTRACT

K. NAKANO, A. KAMIYA, M. IWATA" K. OSHIMA'

Govemmentallndustrial Research Institute Nagoya 1-1 Hirate-cho, kita-ku - NAGOYA - Japan

*Noritake Co. Ltd Miyoshi-cho, Aichi-gun Aichi-pref., Japan

Fiber reinforced silicon carbide, silicon nitride, and mullite composites, derived from continuous fiber, have been developed for appl­ications requiring extreme toughness. They consist of unidirectional carbon fiber reinforcing matrix of the composites. The composites are fabricated by a process consisting of slurry impregnation followed by hot pressing. Measurement of sintering characteristics, identification of crystal phase and observation of metallographic structure in the composites are carried out. Flexural strength and fracture toughness of the composites are measured.

I NTRODUCTI ON

Ceramics have high strength at high temperatures and excellent heat resistance, whereas they have low toughness and reliability for application of structural materials. Several methods have been trying to improve the toughness of ceramics. Fiber reinforcement, as well as whiskers, is one of hopeful method to toughen the ceramics. Several pro­cesses have been employed for fabricating fiber reinforced ceramics. One example is CVI/1-4/, which allowes depositing of vapor reactant of cera­mic material in the interstices of a preform consisting of a fiber ske­leton. Another example is the slurry method/l, 5-7/, which allowes the formation of a ceramic matrix inside of the preform by sllury impregna~ tion followed by sintering either in ordinary pressure or hot pressing.

The CVI method is costly in practice, because it has a long proce­ssing time. However, its reaction temperatures are relatively low, so damage to the fibers are not severe. The sllury method has a short proce­ssing time, but the sintering temperatures are high, so damage to the fiber is unavoidable to some extent. In the present work the carbon fiber reinforced silicon carbide, silicon nitride, and mullite composites were

382

fabricated by the slurry process and examined their characteristics.

I - EXPERIMENTAL PROCEDURE

1.1. Fabrication of ceramic composites

Two continuous carbon fiber groups the PAN group(Tohobeslon Besfight HM-40 and IM-400: 7 ~m~; the structures were amorphous-like graphite), and the pitch group(PETOCA HM-50: 10 ~m~; the crystal structure was crystalline graphite) were used in this process. Fine powders of S-SiC [Ibiceram containing 9 % of A1B2 as a sintering additive(SA): 0.3 ~m], a-Si N (Ube C-OA containing 5 % Y 0 and 5 % Al 0 as a SA: 0.6 ~m) and mullft~(0.7 '~m; stoichiometric co~pdsition) wer~ dsed as fillers. An unidirectional fiber aligned prepreg was made by the filament winding method in which a slurry was used. The slurry consisted of filler, a organosilicon resin[either polysilastyrene(Nihon Soda PSS-400) or poly­silazane(Chisso NCP-200)] and a toluene solvent. The prepreg was cut into segments of equal length[25 mm width and 50 mm length(fiber direc­tion)] to form post forms. The postformes were isostatically pressed and then pyrolyzed up to 700°C in Ar atmosphere. Composite ceramics were made by the sintering(hot pressing) of pyrolyzed postforms at 1600 -1800°C in the Ar. A composite sample fabricated by the aforementioned process is called unidirectional fiber reinforced composite. The process variables used in making the composites are detailed in Table 1.

1.2. Characterization of the composites

The sintering characteristics were measured in regard to open poro­sity and the apparent and bulk densities, while the identification of crystal phases was done by powder X-ray diffraction, and the meta11ogra­phic structure was observed by examining EPMA microgrphs. The flexural strength[room temperature(RT) and 1200°C(in air)], and the fracture tsu-9hness(K ) were measured examining a sample whose size was 3x4x40 mm (the loaacwas applied perpendicular to the fiber axis; the flexural strength: 3 points bending; Klc : Single edge notched beam method).

II - RESULTS AND DISCUSSION

2.1. Characterization of the composites

The sintering characteristics of the composites are shown in Table 2. I n genera 1, the open poros i ty dec reased and the appa rent and bulk densities increased with the hot pressing temperature. In the case of the silicon nitride composites, by comparing the values of the sintering characteristics of the samples using polysilastyrene resin(sample No. 5 -:'8) with those of samples using the polysilazane resin(No. 9 - 12) while holding the fiber and hot pressing temperature variables constant, one can see that the corresponding table values are not so different each other. However in the case of mullite composites, by comparing the values of the sintering characteristics of the samples using polysila­styrene(No. 15 and 16) with those of the samples using polysi1azane(No. 17 - 20), the former samples have higher open porosities and lower appa­rent and bulk densities than those of the latter ones.

The identification of crystal phases in the composites detected by

383

X-ray diffractlon is shown in Table 3. In the case of the silicon carbri ide composites, when PAN fiber was used a S~SiC crystal phase and an amorphaus-like graphite phase were detected(No. 1 and 2), but when pitch fiber was used a S-SiC crystal phase and a crystalline graphite phase was detected(No. 3 and 4).

In the case of silicon nitride composites fabricated by using poly­silastyrene, the samples incorporated PAN fiber produced an amorphous -like graphite phase(No. 5, 6) while the samples using the pitch fiber produced a crystalline graphite phase(No. 7, S). However when the comp­osites were fabricated by using polysilazane the samples using PAN fiber produced a Si?ON? phase in addition to an amorphous-like graphite(No. 9 - 10) phase and the samples using pitch fiber also produced a Si ON 2 phase in addition to a crystalline graphite phase(No. 11, 12). [~ll 4 sample types produced ~- and S-Si N phases]. A hypothesis explaining the presence of Si?ON? crystals i~ that the polysilazane absorbes mois­ture in the air, tnereby forming Si 20N 2 crystals after pyrolyzation.

In the case of mullite composites, by using polysilastyrene and hot pressing at 1600°C, 3Al?032SiO? crystals, amorphous-like or crystalline graphites and SiC phase~ were oetected(No. 13, 15), while hot pressing at 1700°C, in addition to these crystal phases, Al 0 crystals were detected(No. 14). The evidence of the latter case(~ot pressing at 1700 t) may suggest that the mullite decomposes, and A1 20 is formed in the presence of carbon fiber around l700·C. However, by tlsing polysilazane, a considerable Sialon phase(Si1.2AllS01qNs) was detected in addition to the ~- and S-Si N crystals anQ th~ amorphous-like or crystalline gra­phite phases(No~ is - 20). It can be presumed that the pyrolyzation of the moisture absorbing polysilazane forms Si 1N4, Si 20N? and SiC. Evide­ntly either Si N4 or Si ON reacts with the murlite filler forming the Sialon phase. in the ca~e 6f using polysilazane, PAN fiber and hot pre­ssing at l700·C(No. lS), an appreciable Al?01 phase was detected as was the case with polysilastyrene use(No. 14) Onoer same conditions.

Figs. 1 and 2 show EPMA micrographs of planes perpendicular to the fiber axis(the samples represented by these micrographs were silicon nitride composites fabricated by using polysilazane and hot pressing at 1700°C). In Figs. l-(a)[PAN fiber] and 2-(a)[pitch fiber], showing,back scat~eFed electron images, the adhesions between fibers and matrix were tight in both fiber cases. No cracks~ which would presumably arise from a thermal expansion mismatch, were observed around fibers. The fiber volume fraction(Vf) of these samples was estimated to be around 35 %. Figs. l-(b) and 2-(b) show the Si K~ characteristic X-ray images corres­ponding to the micrographs of Figs. l-(a) and 2-(a), respectively. In the micrographs one can observe that Si atoms diffuse far inside the PAN fibers[Fig. l-(b)], while they diffuse only on the periphery of the pitch fibers[Fig. 2-(b)]. This may be attributed that the pitch fibers are crystalline state but the PAN fibers are amorphous-like one, conse­quently, Si atoms can easily diffuse into the PAN than into the pitch fibers. Many microcracks were observed(by SEM micrograph) on the surface of the PAN fibers heated at 1700°C in Ar, while they were not on the surfaces of the pitch fibers. This is another reason that Si atoms 'can easily diffuse into the PAN fibers. Similar observations described 'above can be obtained in the silicon nitride composites fabricated by using polysilastyrene.

Fig. 3 shows EPMA micrographs of the mullite composite, by using the PAN fiber and polysilazane and hot pressing at 1700 0 C. In Fig. 3-(a),

384

showing back scattered electron image, the adhesions between fibers and matrix were tight. No cracks, which would presumably arise from the thermal expansion mismatch, were observed around fibers. The Vf of the mullite composites was estimated to.be around 35 %. Fig. 3-(b) shows the Si Ka characteristic X-ray image corresponding to the micrograph of Fig.~'3-(a). In this figure, one can observe that Si atoms diffuse far inside of the PAN fibers as is the case of silicon nitride[Fig. l-(b)]. Similar image to Fig. 3-(b) can be obtained with Al Ka X-ray beam. In the mullite composite by using pitch fiber and polysilazane and hot pressing at 1700°C, the tight adhesion between fibers and matrix and no cracks around fibers were observed by the back scattered electron image. In the micrographs taken by the characteristic X-ray beams, it can be seen that the Si and Al atoms diffuse only on the periphery of the pitch fibers as was the case with the silicon nitride composite[Fig. 2-(b)]. Mullite composites fabricated by using polysjlastyrene produced mettall­ographic structures similar to the composites fabricated by using poly­silazane.

2.2. Mechanical properties of the composites

Table 4 shows the flexural strength and the fracture toughness of the composites. In the case of silicon carbide composites, by comparing each corresponding fiber type, it can be seen that the flexural strength of the samples hot pressed at 17000 C was higher than those hot pressed at 16000 C(No. 2 >No. 1, No.4 >No. 3). In the case of fracture toughness , by comparing the K values in the samples using the same fiber type and resin type, one t~n see that the corresponding table values are little affected by the sintering temperature.

In the case of silicon nitride composites employing the same resin type and the same PAN fiber type, the room temperature strength and fracture toughness were higher (or nearly equal) when hot pressed at 1600°C than when hot pressed at 1700°C(No. 5 >No. 6, No.9 >No. 10). However, in the case of the samples fabricated by using pitch fiber, an opposite relationship was obtained(No. 8 >No. 7, No. 12 > No. 11). When using polysilastyrene, wHhin each fiber type the flexural strength of samples tested at high temperature tended to have lower values than these samples tested at room temperature. However, the samples using polysilaaane had the opposite tendency. The reason for these results has not been discerned yet, however, it might relate to the fact that the samples fabricated by using polysilazane had a Si?ON2 phase.

In the case of the mullite composites, when uSing the same resin type and fiber type, both the room temperature strength and fracture toughness of the samples fabricated by hot pressing at 1700°C tended to have higher values than the samples fabricated by hot pressing at 1600 t. Both the room and high temperature strengths of the samples fabrica­ted by using polysilazane(No. 14 - 20) tended to have higher values than those fabricated by using polysilastyrene(No. 13 - 16). The reason for these results has not yet been discerned, however, it might relate to the fact that the former samples had a Sialon phase(Table 3).

REFERENCES 1 - E. Fitzer and R. Gadow, Am. Cer. Soc. Bull., 65 (1986) 326 - 335 2 - A. J. Caputo et a1., Am. Cer. Soc. Bull., 66 (1987) 368 - 372

3 - P. J. Lamiq et al., Am. Cer. Soc. Bull., 65 (T986) 336 - 338 4 - H. Hannache, J. M. Quenisset and R. Naslain, J. Matr. Sci., 19

(t984) 202 - 212

385

5 - K. M. Prewo, J. J. Brennan, and G. K. Layden, Am. Cer. Soc. Bull., 65 (1986) 305 - 322

6 - J. K. Guo et al., J. Mater. Sci., 17 (1982) 3611 - 3616 7 - K. Nakano et al., to be published in the "Proc. 87 Int. Symp. Sci

Tech. Sintering" (T(l)kyo, Japan, 1987)

Table 1 Fabrication conditions 01 the cerllic co.posltes

SalPle Matrix of cOIPolte Fiber R •• ln Filler

NIl

1 Silicon carbide PAN (HUO) Po Irsi laatrrene /I-SIC

2 1/ 1/ 1/ 1/

3 1/ pitch 1/ 1/

4 1/ 1/ 1/ 1/

5 SIlicon nitride PAN(lUOO) II a-ShN4

6 1/ II II II

7 II pi tch II 1/

8 II II II II

9 1/ PAN(JM4oo) Polrsllazane II

10 1/ II II II

II II pi tch II II

12 1/ 1/ II II

13 ,ui lite PAN(lUoo) Po I rs Ilashrene 3AI.0,2S10.

14 1/ 1/ II

16 1/ pitch 1/

16 II 1/ II

17 1/ PAN( IUoo) Polrall.zane

16 1/ 1/ /I

19 /I pitch /I

20 II /I /I

Table Z Sinter in. characterlstles of the cerlilc co.posltes

S •• ple li1. 1 2 3 4 5 6

Open 11.6 10.1 B.3 8.B 11.3 6.2

porosltr (") Apparent dens I tv

2.43 2.62 2.69 2.76 2.34 2.60 (./00)

Bulk denslh 2.16 2.27 2.37 2.68 2.10 2.37

(./cc)

S .. ple NIl II 12 13 14 16 16

Open 11.0 6.7 16.7 7.6 porosi t), (")

Apparent deRail)' 2.46 2.70 2.78 2.83

(./c.)

Bulk densllY 2.27 2.66 2.31 2.36

(./c.)

S.IPle nUlber in this table corresponds to that of Table 1.

An error in each value Is .lthln 31.

II

1/

1/

II

/I

/I

/I

7

11.6

2.40

2.16

17

2.2

2.64

2.48

Hot pressin, te.p. ("C)

1750

1800

1750

1800

1600

1700

1600

1700

1600

1700

1600

1700

1600

1700

1600

1700

1600

1700

1600

1700

8 9 10

6.3 10.6 6.1

2.66 2.36 2.53

2.67 2.16 2.41

18 19 20

2.6 2.8 2.4

2.67 2.60 2.62

2.60 2.47 2.66

386

Table 3 Crvstal phases In the cernlc cOIPosites detected bv po.der X-rav diffraction

Filler Fiber Resin Detected crvstal phases

S8Iple No. I n the COlPOS ites

PAN(HII40) Polrsllastnene {J -SIC. C 1, 2 {J -SIC

pitch /I {J -SIC. C' 3, 4

PAN((1I400) /I a-. fJ -SI.N •• C 5. 6

pitch /I {J -. a-SI.N •• C' 7, 8 a -SI.N.

PAN((1I400) Polvsllazane {J -, a -SI.N •• SI.ON •• C 9, 10

pitch /I C', {J -, a -SI.N •• SI.ON. 11. 12

PAN((1I400) Pobsllastnene 3AIo0.2SI0 •• C 13 3A 1.0.25 10 •• AI.O •• SIC. C 14

pitch /I 3AI.0.2SI0 •• C'. SIC 16 3A 1.0.25 10.

3AI.0.2SIO •• Sialon", AI.O •• {J-SI.N. PAN((1I400) Po ball azane C

18

pitch /I 3AI.0.2SI0 •• Sialon". C' 19 3AI.0.2SI0 •• Sialon". C' 20

C araphite (broad pattern: I.e. like norphous phase)

C' graphite (sharp pattern: i.e. like crrstal phase)

# Si .. AI .. O •• N. (ASTII Card NO.31-32)

Table 4 Flexural stren,th and fracture tou,hness of the cerallc cOIPosltes

SalPle t-iI. 1 2 3 4 5 6 7 8 8 10

Flexural RT 29B 418 2Bl 330 294 287 351 6B6 331 190 Strength

(IIPa) 1200 "C 214 410 294 443 382 83

K,e RT 16.B IB.4 10.0 9.B 23.0 12.2 24.6 28.9 22.4 6.8 (IIPa{'l)

S8Iple 11 12 13 14 16 IB 17 18 19 20

Flexural RT 405 686 421 466 406 466 831 497 620 698 Strength

(IIPa) 1200 "C 807 789 483 210 428 603 811 91 728 866

K,e RT 26.7 28.8 IB.4 19.3 14.6 16.1 IB.B 19.3 16.3 17.7 (IIPa{'l)

SalPle nUlber in this table corresponds to that of Table 1.

An error In each value Is .ithln 10%.

387

':'. :.;,. b >, •

Fig. 1 - EPMA micrographs of the silicon nitride composite fabricated by using PAN fiber and polysilazane resin (a): back scattered electron image; (b): Si Ka X-ray image

Fig. 2 - EPMA micrographs of the silicon nitride composite fabricated by using pitch fiber and polysilazane resin (a): back scattered electron image; (b): Si Ka X-ray image

Fig. 3 - EPMA micrographs of the mullite composite fabricated by using PAN fiber and polysilazane resin (a): back scattered electron image; (b): Si Ka X-ray image

ABSTRACT

PROCESSING PARAMETERS INFLUENCE ON THE MORPHOLOGY AND MECHANICAL

PROPERTIES OF SHEET MOULDING COMPOUNDS

J. IMAl, A. RUBIO, C. FORURIA, J.F. LlCEAGA

Inasmet BO Igara sis 0 20009 SAN SEBASTIAN 0 Spain

The weight reduction due to the utilization of the SoM.C. makes this material appropiate for several applications. The final property of a S.M.C. product depends on the process employed for the transformation. Therefore it shall be interesting to know how the processing parameters affect the mechanical and thermal properties of the product. This work intends to study the effect of mold pressure mold temperature and cure time on the flexural properties and on the Tg of the material.

INTRODUCTION

The S.M.C., due to its low cost, good mechanical properties and low times of production, is being introduced for various applications in different industrial sectors. One of the most important characteristics of the work with this kind of materials is the strong dependence of the processing parameters on their properties. There are different publications about the relationship between the process parameters and the mechanical properties of the materialso/l/,/2/. However there is little bibliography joining these two factors to the material morphology and its Tg.

In this paper the influence of the process parameters (mold temperature, cure time and mold pressure) is studied over the flexural properties of a S.M.C. and its relationship with the thermal and structural properties. The Tg has been studied by thermal mechanical analys (T.M.A.) methods.

MATERIALS AND EXPERIMENTAL TECHNIQUES

A chopped glass fibre reinforced ortophtalic polyester S.M.C. has

390

been used for the experimental work. The composition of this S.M.C. is 35% polyester resin, 25% glass fibre and 40% hydrated alumina filler.

The S.M.C. has been processed in a 100 T press. Both cure time and mold temperature have been studied and modified processing parameters. The variation interval on the cure time is between 2 and 7 minutes, on the pressure between 2 and 8 MPa and on the temperature between 130 and 1809C.

After the transformation the morphology of the material has been characterized through Tg measurements. The Tg has been measured in a Thermomechanical Analyser, observing the change in the specific volume of the material against the temperature.

The flexural test were carried in an Instron 6025 Universal Testing Machine following the ASTM D-790 standard. Using the flexural tests, both fracture strength and flexural modulus are obtained. The variation of these mechanical properties with the processing parameters is also studied. The specimens were conditioned at 249c and 50 vercent relative humidity for a minimum of 48h before testing.

RESULTS AND DISCUSSION

The evolution of glass transition temperature Tg respect to the cure time may be observed in fig. 1. The Tg increases with the time possibly due to the reduction of the mobility in the polymeric chains due to a better curing of the resin.

Fig. 2 shows the influence of the mold temperature on the Tg of the product. The increase of the temperature produces a higher curing rate and therefore a higher Tg.

In fig. 3 it can be also observed a little increase of the Tg as a result of the mold pressure raise, although this increment is qualitatively lower than the previous case. The pressure may reduce the microvoids which may reduce the chain mobility and increase the Tg.

Fig. 4 and 5 represent the increase of the flexural strength and the modulus when mold temperature raises. This increment on the mechanical properties is due to the better curing rate shown by the Tg values obtained in fig. 2.

In fig. 6 and 7 the increase of the flexural strength and modulus is observed when the mold pressure is raised. This increase of the mechanical properties is due to a reduction of the void volume in the piece.

Finally, figures 8 and 9 represent the effect of cure time. Both parameters, flexural strength and modulus show a maximum value for intermedium curing times. Lower values obtained for short times

391

may be due to the undercuring of the resin and low values shown for long times to an overcuring of it, producing some kind of degradation. This behaviour is in concordance with some similar effects studied in the bibliography./l/, 13/.

CONCLUSION

The influence of some processing parameters of a S.M.C.(mold temperature, mold pressure and cure time) on the flexural pro~erties and Tg has been studied. A variation on the Tg has been observed with the three parameters and it is more sensitive in the case of the temperature and time due to a higher cure degree.

These effects can be observed also on the mechanical properties. In this way it is possible to adjust the process parameters for the optimization of the material properties.

100,-----------------------------------,

_.__----a--u a ______ ---------

80

0< 70 ______ --

" E-< 60

2 5

TIME (MIN)

Fig 1 Influence of cure time on Tg

7

l00r---------------------------------------,

80

u 70 C>

TEMPERATURE (QC)

Fig 2 Influence of mold temperature on Tg

8

392

100

90

u 80 ()

l!) 70 Eo<

60

50 0 2 4 6 B 10

PRESSURE (MPa)

Fig 3 Influence of mold pressure on Tg

125

-;;100 ~ ~ 7S z

~ '" 50

~ 25 ~

0 120 130 140 150 160 170 180 190

TEMPERATURE (QC)

Fig 4 Influence of mold temperature on flexural strength (t=3 min and P=4 MPa )

nxor---------------------------------------, -;; ~ 6000

I i SOOO

~ 4000 ~

TEMPERAWRE (Q C)

Fig 5 Influence of mold temperature on flexural modulus (t=3min and P=4 MPa)

393

125 -;0

'" ~ 100

I '" Eo<

[;1 75

~

! Eo<

! t 0 III

~ :50

" >: OJ 25 H ..

0 0 2 4 6 8 10

PRESSURE (MPa)

Fig 6 Influence of mold pressure on flexural strength (Temp.=1309C and t=3 min)

6000

~ 5500

"'

f ? " 5000 §l I §l

~ f -' 4500 ;:i " >< '" -' 4000 ...

3500 0 2 3 5 6 7 8 9 10

PRESSURE (MPa)

Fig 7 Influence of mold pressure on flexural modulus (Temp.=1309C and t=3 min)

394

100

.. 75 !l § 2: 50 ~

I I III ,..,

~ 25

'" Ii

0 1 2 3 " 5 6 7 8

rIME ( MIN.)

Fig 8 Influence of cure time on flexural strength (Temp.=1509C and P=4 MPa)

7000

.. !l 6000 ~ III

I i 5000 ~ ,..,

f ~ " >< '" 4000 ,.., ..

3000 1 2 3 4 5 6 7 8

lIME (HIN.)

Fig 9 Influence of cure time on flexural modulus (Temp.=1509C and P=4 MPa)

REFERENCES

1.Tung R.W. "Effect of processing variables on the mechanical and thermal properties of S.M.C." ASTM. STP 772 pp 50-63

2.Mallick P.K. & Raghupathi N. Polymer Engineering and Science. Vol 19. N911. (1979) P 774

3.Morton M. "Rubber Technology" 2nd ed. Van Nostrand Reinhold. New York. Chapter 4.(1973)

ABSTRACT

THE MANUFACTURE OF ULTRA-LIGHTWEIGHT LARGE DIAMETER COMPOSITE PISTONS

P. MOBBS

Royal Aerospace Establishment Materials and Structures Department X34 building

GU14 6TD FARNBOROUGH HANTS - England

Heat transfer characteristics of gas turbine blades can be measured in an intermittent short steady-pulse wind tunnel. Such a tunnel has been constructed at RAE Farnborough in which a stable hot gas pulse is generated by an ultra-lightweight, large diameter piston. Manufacture of such a piston is made possible by composites fabrication techniques recently developed for aerospace applications. Two pistons have been fabricated and found to exceed their design and operational requirements.

1 INTRODUCTION

Improvements in the efficiency and performance of gas turbines can be achieved by increasing operating temperatures. Accurate assessment of heat transfer to turbine vanes is therefore required to permit optimisation of. blade cooling systems. A facility has been developed by Oxford University and RAE Propulsion Department to allow heat transfer and aerodynamic effects to be studied under conditions accurately simulating an engine environment. The facility comprises a novel form of wind tunnel which provides a short duration, steady flow through a cascade. A schematic layout of the test rig known as the Isentropic Light Piston Annular Cascade facility1 is shown at Fig 1-

In a typical test, high pressure air is admitted to the

396

cylinder behind the piston, driving it forward to compress and heat the mass of air in the cylinder ahead of the piston. When a predetermined combination of pressure and temperature is achieved the fast acting plug valve is opened allowing the heated air mass to flow through the test cascade where heat transfer into blades or vanes can be measured. Two such rigs have been built, a O.6m dia. tunnel at Oxford University, and a 1.2m dia. tunnel at RAE Farnborough. This paper outlines the design features and subsequent fabrication by RAE Materials and structures Department of ultra-lightweight composite pistons for the two facilities.

2 DESIGN CONSIDERATIONS

Both test rigs operate at pressures up to 16 bar, this value currently being limited by piston performance. The mass of the piston is a critical factor; it must be less than the mass of air in the pump tube in order to avoid unacceptable pressure oscillations. This gives a nominal mass of 4.5kg for the O.6m dia. piston and 24kg for the 1.2m dia piston. However, a high structural efficiency is required. For example, the 1.2m dia. piston must be capable of transmitting energy pulses of up to 10MJ/stroke at a 10MW power rating.

Initially, a piston for the O.6m dia. tunnel was fabricated at Oxford University from aluminium alloy sheet and using conventional fabrication methods. It was not possible to obtain sheet material of optimum thickness for the design and due to the expense and time that would have been required to machine the sheet to optimum thickness the piston mass was excessive. This caused unacceptable unsteady gas flow conditions in the tunnel.

At this stage it was decided that future pistons should be manufactured using composite materials both to save weight and to reduce fabrication costs. The use of sandwich materials allows the fabrication of lightweight structures having exceptionally high efficiency when subjected to bending loads. Also, the ability to fabricate sandwich skins of any thickness by selecting the number of plies eliminates the need for the expensive machining and surface preparation of the metal structure. For construction of these pistons, fibre composites therefore have the potential to reduce both weight and fabrication costs.

A prototype piston was fabricated for the 1.2m tunnel by Oxford University,2 to operate at pressures up to 8 bar This was made using commercially available flat panels of CFC skinned aluminium honeycomb core. The piston met the required weight target and initially performed well. However, after a limited number of test cycles the piston crown ruptured causing extensive damage to the tunnel cascade section. It was then decided to fabricate future pistons in-house, at the R.A.E composites Structures

397

research facility. This would offer a potential to upgrade the piston performance by optimising the combination of sandwich core density and skin lay-up to resist the specific loads experienced by the structure.

3 PISTON MK 1 (1.2m DIA 16 BAR.)

3.1 Design concept

The schematic layout of the larger piston is shown at Fig 2. It consists of a honeycomb sandwich crown with a moulded GRP flange to locate the piston ring and a circular skirt. In order to minimise requirements for machining and fitting operations and to reduce the parts count, the maximum possible potential for moulding to the finished shape was sought. In view of the one-off nature of the project, tooling costs had to be constrained to an acceptable level and therefore, using a modification of the guidlines described by Lynch 3 , low-cost wood and rolled aluminium plate tooling was used.

Ciba-Geigy . XASj914 pre-preg. was chosen for the manufacture of all CFC components because it has tolerant processing characteristics, because considerable experience in using this material had been gained at RAE and because a large data bank of material properties existed. Design calculations were performed and coupon specimens were fabricated and tested in order to verify the properties of the combination of honeycomb core and composite skin chosen in the design.

3,2 Design specification

The design specifications were as follows:

piston crown core; Hexcell type 5052, 75mm thick, cell size 0.125in, aluminium alloy gauge 0,002in.

CFC skins; Ciba-Geigy Fibredux XAS/914 pre-preg. 4 plies, lay up 0,90,90,0.

Glass fibre cloth on front face to improve tolerance to impact.

Bonding carried out using Ciba-Geigy BSL312/5, 120°C film adhesive.

Piston ring groove shell; glass fibre cloth, wet lay-up moulded using Ciba-Geigy resin type LY 568

Piston skirt core; Ciba-Geigy Aeroweb AI-64-3 Nomex honeycomb, 10mm thick, cell size 3mm.

Skirt skins; Fibredux XASj914 pre-preg. 3 ply lay-up 0,90,0.

3.3 Fabrication

Fabrication of the crown was carried out in four operations. Initially the top and bottom skins were laid up in the form of circular disks and cured. The GRP piston

398

ring housing was layed-up on wooden tooling as a separate operation and subjected to a hot cure cycle using methodology developed by Childs4 • The honeycomb core was trimmed to a disk shape using a bandsaw. The skins wer,e bonded to the core and piston ring housing using film adhesive in a one shot autoclave cure cycle.

The piston skirt was fabricated in two operations. First the inner and outer skins were layed-up and cured and then bonded on to the honeycomb core material using film adhesive. 'Springback' occurred during this operation and had to be allowed for by trimming the two half skirt sections to butt together forming as near a circular section as possible. All components were subjected to ultrasonic quality assurance tests before and after assembly.

Final assembly involved bonding the two half-skirt sections to the crown with butt straps at all joints as shown Fig 2. Prior to commissioning trials a protective two part epoxy paint finish was applied.

4 PISTON MK 2 (0.6M DIA. 16 BAR)

4.1 Design concept

As a direct result of experience gained manufacture and operation of the first piston design modifications were made for the smaller schematic layout is shown at fig 3.

during the a number of piston. A

Difficulty was encountered whilst maChining the honeycomb core material for the first piston. To eliminate this problem, a rigid foam was used as the core material for the second piston. Since stiffness was not a critical parameter, it was decided to use a woven Aramid 120°C cure pre-preg. This material has a high specific tensile strength, exceptionally good impact resistance and the use of woven fabric pre-pregs allows the fabrication of thin section isotropic plates. A high resin content system was selected to enable curing and bonding operations to be carried out simultaneously without the addition of film adhesives. A further design change was to use a square skirt. This would eliminate dimensional errors arising due to 'springback' and, in addition, the square structure would be more stable than a circular skirt when subjected to the side loads from the loading pads.

4.2 Design specifications

Crown core; Rohacell WF 200, 25mm thick. Crown skins; 2 plies of Kevlar cloth, Ciba-Geigy Fibredux

920K-285-52% Skirt core; Rohacell rigid foam, Grade WF 110 Skirt skins; 2 plies of Kevlar, Ciba-Geigy Fibredux

'920K-285-52% Lightweight filler Ciba-Geigy Redux 252

399

4.3 Fabrication

Initially the foam core was machined to the correct diameter and the pre-preg. blanks were trimmed to their final size. These components were assembled and then cured

and bonded in a single autoclave cure cycle. It was also possible to fabricate the skirt in a similar one-shot cure operation. The skirt was then bonded to the crown using Kevlar pre-preg. to form the butt straps. Final manufacturing operations required filling the crown periphery with a lightweight syntactic foam to form a machinable layer into which the piston ring groove could be cut and bonding on the 4 corner support pads. A protective two-part epoxy paint was applied prior to commissioning trials.

5 CONCLUSIONS

Both pistons have passed acceptance trials. The 1.2m diameter piston has now been operating for a period of 4 years completing over 750 cycles in the tunnel facility with no detectable damage. The programme has demonstratedthat it is possible to fabricate one-off high performance structures from composites at competitive cost. High material costs and autoclave running costs are offset by the use of low cost tooling, and the use of net curing which eliminates a much of the expensive machining and material waste that would be associated with alternative fabrication routes. It is probable that the construction of a piston having acceptable strength/weight characteristics could not be achieved without the use of high performance composites.

REFERENCES

1 Brooks A.J. The Isentropic Light Piston Cascade Facility. RAE Technical Report 85098 (1985)

2 Brooks A.J., Colbourne D.E., et al. The Isentropic Light Piston Cascade Facility at RAE pyestock. RAE Technical Memorandum P 1053 (1985)

3 Lynch S. Characteristics and Performance of Tooling Fixtures for Curing Fibre Reinforced Thermosetting composites. RAE Technical Memorandum Mat/Str 10091 (1987)

4 Purslow D. Childs R. Autoclave Moulding Of Carbon Fibre Reinforced Plastics. RAE Technical Report 85039 (1985)

copyright (C) Controller HMSO, London, 1988

400

LIGHT PISTON

Fig.1 - Schematic layout of the test facility.

GlASS ~I8AE TOP HAt AJU

PTF'E PIS TON AING

PlSTONAIlooG~

EPOXY'uEA

.,""'" E':;;S;!---~UUINlUU H()t<.EYCOue

GlASS FIBRE SCRllol

CAABQNfllIBRE: fACES

CUAI!. OF SU"O'-T !"ADS

CArtDCft l llllVlr 'AGrn ,1()N{'fC(:II,IIH 91.'"'

c rAP Bun JOINT

Fig.2 - The 1.2m dia.piston

OE:TAIL ~ IDR"£A REI~CA:IN3.

OETAIl cr PISTCN AIN(l MOUSIN3.

DETAIL OF JOINT REINFORCING.

Fig.3 - The O.6m dia. piston

ABSTRACT

DESIGN OF DOMES BY USE OF THE FILAMENT WINDING TECHNIQUE

M. MARCHEDI, D. CUTOLO', G. DI VITA

Aerospace Department University La Sapienza via Eudossiana n° 16 - ROME -Italy

'SNlA Colleferro - ROME - Italy

This paper will show a state of the art idea in the sector dealing with the design of domes that are accomplished in Filament Winding (F_W.). Later they can be used for either a pressure reservoir (only membranal loads) or as a covering for the rocket's motor (membranal and concentr~

ted loads). While undertaking the study of winding in the above cases we

used two different hypothesis: without friction between the fibers and

the underlying surfaces as well as conditions with friction. A computer aided winding system of composite structure has been developed. Such ap­proaches opportunely adapted, can be utilized to determine the superior

winding method of an assigned aXisymmetric design.

IN'l'RODUCTION

In the winding process the most important elements are the mandrel configuration, the fibers orientation, their trajectory on the mandrel surface and the movement of the eye of the winding machine [1), [2). For

what concerns the trajectory of the filament, it is well known that, for every curved surface, the greatest winding stability is obtained follow­ing the geodesic lines. This technique is limited to the simplex compo­

nents; for complex shapes it is necessary to use the non-geodesic traje~

tories [3), [4). In this paper you find the philosophy of coverings rea­

lized in F.W.: aXisymmetric structures generally formed by a central cy­

lindrical body and by two domes utilizing helicoidal and circumferential

winding. You'll find here studied the domes for reservoirs exposed to di!

ferent loads with different dimensional requisites, estimating the cons~ quences in terms of the design parameters. You also find suitable models to simulate windings in the CAD/CAM systems [2), [4) for any axisymmetric geometry .

402

1. PRESSURE VESSELS: BASIC EQUATION

On the basis of the classic theory of pressure vessels [1], the thick

ness of filaments in the cylindrical part (Fig. 1) are:

3 cos 2 a o

ttot (2 - 3 sin 2 a o)

3cos2 a o

total thickness of wall

helicoidal thickness (1 )

circumferential thickness.

As far as the dome is concerned, the thickness of the helicoidal winding is as follows:

(2)

The configuration of both dome and filament trajectory is determined by supposing the contribution of resin to be negligible (net-method) and the stress a to be the same in all fibers and constant for all their length (isotensoid design). Knowing the loads, by the translation equilibrium of filaments we obtain the expression of the winding angle a:

2 re tga=2-­

r",

(3)

and also, by the translation equilibrium and the thickness variation (2), the stress in the filament:

pr2 a = -----=---7""'"--

2 ro teo cos ao cos B cos a

which is constant if:

cosa cos B = constant (4)

The connections between r e, r"" rand B allow us to express the (3) and (4) as follows:

rr" tg2a 2 + (r') 2 +

r2[1 + (r') 2]1/2 (5)

= k cosa

consisting in a final system for the determination of rand a(z). By sim plifying eq. (5) we obtain the Clairaut's equation:

r sina. = constant (6)

which defines the geodesic lines on a revolution surface.

403

We obtain also a relationship between the nondimensional height Z z/ro and the radius X r/ro:

dZ

dX (7)

which allows to draw the dome meridian. For the study of the trajectories we use a differential equation in order to determine the angle e when the dome geometry and the winding law a(z) are known:

de dz

11+(r,)2

r tga

2. WINDING WITH CONSTANT SLIPPAGE TENDENCY

(8)

A non geodesic winding is stable and suitable if the slippage tende~ cy of filaments is not greater than the slippage resistance caused by the friction. A satisfactory stability control can be obtained making the slippage tendency constant in every point. By indicating with ~ the angle formed by the main normal it of the line r and by the normal rts to the S surface, Fig. 2, it is evident that on each uni tary element of fiber a lateral -Tc sin ~ bs and normal Tc cos ~ rts for­ce develop. In order to avoid the slippage it must be:

We define "slippage tendency" k the ratio between the lateral and the nor mal force:

Tc sin ~ k =

Tc cos ~ = tg~ (9)

The condition for winding stability, along the non-geodesic path, can be expressed in the following manner:

Itg~1 .;; tgtp = jJ

where tp is the dome/fiber friction angle. The relationship between ~, a and r is (2):

tg~ = (a'r cosa +r' sinal [1 + (r') 2) rr"cos 2a-[l+(r')2) sin2 a

and then we have the final set of equations:

2 +

k

e' =

r rtl

1 + (r') 2

(a'r cosa+r' sina)[l + (r')2) rr"cos 2 a-[l+(r,)2) sin2 a

[ 1 + (r' ) 2)112 tga

r

(10)

(11 )

(12)

It is interesting to see that if tglP = 0 from the equation (11) we obtain

404

the geodesic equation (6). The Figs. 5, 6, 7 compair the geodesic with

non-geodesic domes with different slippage tendency k = 0.2, k = - 0.2.

3. ROCKETS

Let us consider a typical load system, of solid rocket motor dome, for which internal pressure p and concentrated force F at the dome tip

simultaneously occur (Fig. 3). Also in this case the isotensoid and non­isotensoid designs are considered. In the former case the first equation of (5) is replaced by:

2 r r" 1-N(r o/r)2 + 1+(r,)2

where N is the "non-dimensional load factor":

F

N = p1Tr~

Whereas the second equation of (5) becomes [3]:

r 2/i + (r')2

coso.

(13)

(14)

(15 )

The dome meridian and the winding equation a(r) are determined by solving the following differential system:

2 r '('II

tg 2a 1 - N(ro/r) 2

+ + (r') 2

r2/1 + (r') 2

[1 - N (r: fJ ( 16)

= constant coso.

It is possible to show that, also in this case, the isotensoid design leads to a geodesic winding law. The dome meridian is given by:

dZ dX

(17)

it is evident that for N = 0 we obtain the equation (7). Some numerical results of eq. (17) are shown in Fig. 8. Due to inherent dimensional li­mits of eq. (17) also in this case the non-geodesic (constant slippage tendency design) dome choice is mandatory. It is possible by substitu­ting the eq. (13) in the set of equations (12), [3], to obtain:

2 r r" + = tg 2 a

1-N(ro/r)2 1+(r,)2

k =

8'=

(a'rcosa+r' sina)[l + (r,)2) rr" cos 2 a- [1 + (r')2) sin 2 a

tga r

The relevant result for N = 0.2 and k

( 18)

1.5 is shown in Fig. 9.

405

4. GENERAL AXISYMMETRIC GEOMETRY MANDRELS

The above mentioned theory can be successfully applied for the dete~ mination of the winding paths to be performed on general axisymmetric geo me try mandrels. In case of geodesic windings the path r is given by equ~ tion (6), which r(z) being known, allows to determine a(z). Afterwards, by using equation (8), 8(z) can be obtained. The non-geodesic windings can be studied by replacing the axisymmetric shape with a sequence of truncated cone (Fig. 4). Since for a cone, whose tip angle is 6, the fol lowing equation are drawn:

r' dr

tg6

r" = d~ (~~) = 0

the second eq. of (12) can be modified as follows:

a' = tga (tg6 - ksina) r

(19)

(20)

The problem can be solved on the i-th conical element by integrating the

following system:

da tga (tg 6i - k sinal dz r

de tga (21 )

dz r cos 6

The slippage tendency k, can be varied step by step on the discretised su~ face. The results simulating a conical sample winding are shown in Fig. 10.

5. CONCLUSIONS

In this paper equations have been presented for the design of domes that are accomplished in Filament winding using geodesic and non-geodesic winding trajectories. They permit to develop a computer aided design sof! ware for the filament winding of composite axisymmetric structures. It pr~ vides a basis for determining winding regularity, controlling motion of the fiber feed arm and calculating strength and rigidity of the products. An extension work is in progress to deal with non-axisymmetric shapes and even general tapered structures.

REFERENCES

1. M. Marchetti, D. Cutolo, "Tecnologie dei materiali compositi", ESA­Masson ed., Rama (1987).

2. J.P. Denost, "Design of Filament-Wound Rocket Cases", AGARD n. 150, pp. 5.1-5.21.

3. G. Di Vita, "Thesis", Aeronautical Engineering, Univ. di Roma (1988).

4. G.M. Wells, K.F. Mc Anulty, "Computer Aided Filament Winding Using Non-Geodesic Trajectories", ICCM-VI, London (1987), pp. 1.161-1.173.

406

x a. MERIDIAN

F Il AN E N T (r)

ro: cylinder radius r: dome radius r~: meridian curvature radius re: parallel curvature radius ao: winding angle on the cy-

linder

a: winding angle on the dome Gam: ultimate stress c: curvature

Fig. 1 - A pressure reservoir with its different parameters.

Fig. 2 - Geometry of a non­geodesic path.

Fig. 3 - Rocket's motor scheme.

z z.

y r

Fig. 4 - A general revolution surface.

1.0~-----,.---------,

z a=17.476° o

0.6

0.4

0.2

0.0,4---,---,---.--,.--...-... 0.0 0.2 0.4 0.6 r(z) 1.0

1.0..,.......,---------...,

z

0.6

OA

0.2

ag-17.476°

K=.2

O.U+--L-----r---r---r----,----t 0.0 0.2 0.4 0.6 r(z) 1.0

407

Fig. 5 - Meridian of a geodesic dome with eto = 17.476° (din = 0.3).

o

270r-----~~----~90

180

Fig. 6 - Meridian and up-view of a non-geodesic dome (k = 0.2) with (10=17.476°.

o

t a 0:: 17.476° e

270~~--~--~--~90

0.2 0.4 0.6 r(z) 1.0 180

Fig. 7 - Meridian and up-view of a non-geodesic dome (k = - 0.2) with (10=17.476°.

408

1.0 1.0

Z diD =.2 Z diD =.2

N =.3 0.6. 0.6

0,4 0.4

0,2 0,2

0 0 0 Q2 0.4 0,6 r(z) 1,0 0 0,2 0,4 0,6 r(z) 1,0

Fig, 8 - Meridian of geodesic domes with increasing values of the load factor N.

1,0...-------------.

z

0.6

0,4

02

ao= 11.53°

K = 1,5

N=.2

o'+---~--~--~--~--~ o 02 0,4 0,6 r(z) 1,0

o SPHERIC

270~-i-~~-~~90

180

Fig. 9 - Meridian and up-view of a non-geodesic dome (k=1.Sl with load factor N = 0.2.

-100

I : I I K=.2 I K= 0 ,K=.21S :K= 0 K=-,2

I , I I I , , I

·~)())fr1 I I I Ii: I I , I I I , I

a b -200 L,....--~--,---.-------,-

o 100 200 (mm) 400 o 100 200 (mm) 400

Fig. 10 - Winding simulation on a conic mandrel: geodesic (al and non­geodesic (bl trajectories on the cone.

COMPORTEMENTS MECANIQUES MECHANICAL PROPERTIES

Chairman: Mr Th. JOHANNESSON Institute of Technology of Link6ping

STATISTICAL INFERENCE ABOUT STRESS CONCENTRATIONS IN FIBRE·MATRIX COMPOSITES

L.C. WOLSTENHOOME, RL SMITH, M.G. BADER

University of Surrey· GU/LOFORO . Eng/and

ABSTRACT

It is shown how experimental data on composites and the component fibres may be used to make inferences about the stress concentration factors and length of stress overload region. The method employed is numerical maximum likelihood, involving detailed combinatorial calculations, and is therefore highly computationally intensive. The method is illustrated using experimental data on hybrid composites consisting of carbon fibre tows embedded in glass-epoxy composite, particular emphasis being placed on the consequence of varying the distance between tows.

1-INTRODUCTION

It is generally accepted that in a uniaxial composite loaded monotonically in tension, parallel to the fibres, fibre fractures will occur progressively at the weakest points in individual fibres. At a sufficiently high load the stress concentrations induced in the unbroken fibres close to those fibre breaks will lead to further fibre fractures and eventually to catastrophic failure of the composite. It is difficult to observe individual fibre failures in a composite so this has been modelled with a hybrid system consisting of a uniaxial glass/epoxy laminate which contains an array of carbon fibre tows. The 1000 filament tow has been used to represent a single ligament in a composite. These tows each contain 1000 filaments and the tow is approximately 0.3mm in diameter. In earlier investigations, Bader & Pitkethly [1,21 measured the strength of single carbon-fibre tows impregnated with the epoxy resin and also single tows embedded in the glass/epoxy to form a simple hybrid laminate as described above. In this work 7 carbon-fibre tows have been assembled in the glass/epoxy coupon and the spacing between the tows varied from 1.5 mm (5 tow diameters) downwards, i.e. to tows just touching. The sequence of tow failures was observed as the coupon was extended ( as indicated

412

in Fig.1 ) in order to determine the extent to which the failures in one tow influence those in neighbouring tows. If detailed data are available concerning the positions and failure stresses of individual breaks in the composite, then much can be learned about the stress concentrations. A new method of statistical analysis is proposed which attempts to model the whole process of the appearance of flaws in the material. It is based on the concept that failure in one tow will result in load transfer to the surrounding tows and matrix so that there will be a stress concentration in the tows adjacent to the failure. The lengthwise distribution of stress in this region is controlled by shear interaction between the broken and unbroken tows through the matrix.

The method is based on the principle of maximum likelihood, and requires detailed computations of individual sequences of failures. Numerical estimates of parameters describing the pattern of stress concentrations are yielded but there is considerable uncertainty about these parameter estimates as reflected in the confidence intervals derived. Nevertheless, by this method it is possible to evaluate the differences among stress concentration patterns at different tow spacings, and to make inferences about the lengthwise distribution of stress concentrations in the neighbourhood of a failed tow.

2-STATISTICAL MODELS FOR SINGLE FIBRE STRENGTH

The long established Weibull model for single fibre strength rests on the assumption that failure is due to flaws which occur independently and randomly along the length of the fibre. For the two-parameter Weibull distribution, the probability that a fibre survives stress x is given by

1 - Fa (x) = 1 - exp{-a(x/xl )W}

where a is the length of the fibre, Xl the characteristic stress of the fibre at length 1,

xa = Xl a(-l/w) the characteristic stress at length a,

and w is the Weibull shape parameter.

Suppose a fibre of length a consists of n segments of length d. If all segments can be regarded as independent then

1 - Fa (x) {1 - Fd (x) }"

This is commonly referred to as the "weakest-link" concept, i.e. the assertion that a fibre is only as strong as its weakest portion. The Weibull distribution is consistent with this relation, and is a simple two-parameter function found to be consistent with a wide variety of materials.

Suppose a fibre of length d fails whilst under the influence of transfered load due to the failure of adjacent fibres. This enhancement of effective load may be reflected in the application of a load concentration factor k. When the applied stress is x the fibre experiences stress kx. As failures progress, the load concentration factor increases from its initial value of 1. If a fibre currently surviving stress kx, fails whilst the load concentration factor is increased from k to k-, the probabilty of

413

this event is

Fd ( k-x) - Fd (kx)

i.e. the exact strength of the fibre is unknown, only that it lies between kx and k-x. Suppose, however that a fibre fails at a known experienced stress kx, then the probability of this event is the limit of the above expression as k- -} k which is the differential k fs:J (kx) dx, where f(x) = dF/dx. Such calculations under11e the principle of the method: we are able to calculate probabilities of observed patterns of failed and unfailed fibres, as a function of unknown stress concentration parameters. The method of maximum likelihood then permits us to make inferences about those parameters.

3-MODELS FOR FIBROUS COMPOSITES

Many of the theoretical models for fibrous ~omposites, are concerned with fibres (or tows) arranged 1n a linearly equally-spaced array. The load on any unfailed fibre may be assumed to be a function of how many failed fibres are adjacent to it. If this number is r then k is of the form 1 + g(r). This analysis f~lows Bader and Pitkethly and assumes g(r) to be of the form ~(r)/F, where F is termed the load sharing factor, and reflects the degree to which the load of a failed fibre has been transfered to other fibres.

It is accepted that in practice all sharing fibres do not bear an equal load increment , and that the actual stress concentrations are not as localised as the following model implies, but it is considered that the model represents the essential features of the situation and makes the following calculations possible. In accordance with the "chain of bundles model" (Harlow and Phoenix [3,4]) a set of N parallel tows or fibres (a bundle) is considered as a set of m independent sub-bundles length d. (Fig.2) The length d is intended to represent the 'ineffective length' i.e. the distance in the direction of the composite over which the stress concentrations occur. There is considerable uncertainty as to its true value, and also as to how much it might vary with stress, but it is assumed here to be constant and a number of different values have been tried in the region of 5-10 tow diameters as suggested by the early work of Rosen [5] and Zweben [7]. On dividing a bundle into sub-bundles of length d, it is desirable that breaks which might be related should be in the same sub-bundle, to support the independence notion. Clearly this latter condition cannot be guaranteed and it must be borne in mind that the chain of bundles model is only an approximation.

Given data for the failures occurring in the i'th sub-bundle, L. the jOint probability of the events observed in the N segments making up sub-bundle i may be calculated. Under the given load-sharing rule, appropriate load scaling factors for each segment are calculated, whether it be for the point of failure, an interval in which failure occurred, or when the segment survived the final stress. In addition, where segments apparently failed 'simultaneously', all possible orders of failure must be allowed for , failure presumably triggered by anyone of the group.

The likelihood function for the whole bundle (under independence) is given by L L, L2 L3 L4 ............. Lm•

414

These calculations are sufficiently complex that considerable computer time is involved merely in evaluating the likelihood function which then has to be maximised with respect to the unknown parameters. The strength distribution of individual bundles is required and taken to be adequately represented by the Weibull model. Given the strength distribution of a bundle length L, the strength of a segment length d is deduced from the "weakest-link" relationship. Here L is primarily being considered as a function of F and d , but the Weibull parameters xl and ware also of interest. The strategy that has been adopted is to fix d and one or two of the other parameters and estimate the rest.

The maximum likelihood estimates are those values of the parameters which maximise the likelihood function. In addition, by the standard chi-squared approximation to twice the log likelihood function, approximate confidence intervals may be put on such estimates.

4-RESULTS

For each inter-tow spacing several specimens were used and since they came from the same batch of material, the strength parameters could be assumed the same across the specimens and the data combined to give an overall estimate of F. The Weibull parameters were estimated from first failures of each tow where these could be considered independent of any other failure and also assessed in the light of experiments carried out by Pitkethly on single tows in glass epoxy.

It was expected that F would increase with inter-tow distance, and effectively be infinite at a distance greater than 4 to 5 tow diameters, where it is considered that all the excess load is absorbed by the matrix. Further, the estimate of F is expected to vary with the value chosen for d. If a given load transfer is spread over a wider region the transfer per unit length must be less and therefore would be reflected in a larger value for F. In general terms the results conform to these principles ( Table 1). In the 1.5mm spacing case the confidence intervals for F are finite but very wide. This is to be expected as it is evident that the amount of dependence between breaks is slight, so precise estimates of F are unlikely to be achieved.

The results for the 1.0 mm spaced tows were not very satisfactory. Variations in experimental conditions may be responsible - the estimated characteristic strength for this spacing was considerably lower than usual, and there were significant differences between specimens. Joint estimates of Xl and F were obtained for various values of d.(Table 2). Again the values of F do not vary in a consistent way and the estimated Xl should be independent of d. Closer examination of the data revealed that d = 2 or 4 resulted in far more 'adjacent' breaks per sub-bundle than for the other values. It was surprising that this should happen over a total bundle length of 500mm. It is felt that d = 1.0, 2.5, 5.0 gives the more realistic picture, and within this group the estimated parameters behave in a reasonable way. The specimens with 0.5mm spacing were found to be in keeping both with each other and with the general strength distribution. Estimates of F obtained individually and jOintly with Xl behaved consistently, and estimates of Xl were independent of d. In the case of tows touching, all breaks occurred as part of a break across the entire

415

specimen, the tows behaving as though a single larger bundle (tape). Without any tows surviving neighbouring breaks it is not possible to estimate F.

5-DISCUSSION

The different values of F for different values of d are the result of approximating the stress concentrations around a break by a step function, being k = 1 + g(r) up to a distance d/2 away from the break and thereafter 1. In reality the stress concentration factors in adjacent fibres decay continuously from a maximum opposite the break ( Fig.3 ).' A suitable functional form of this relation is that at a distance t from the break the true stress concentration is 1 + a exp(-ct). Taking r = 1 it was found that the results of the 0.5 Mm. spaced bundles conformed well to this exponential stress decay relation with a = 0.0936 and c = 0.825. Results are not very sensitive to the choice of c but it has been shown that c = 0.825 is acceptable for the other spacings, giving a = 0.0154 for the 1.0mm spaced bundles ( d=1.0,2.5,5.0) and a = 0.013 for the 1.5mm spaced bundles. Further discussion may be found in Wolstenholme and Smith [6] .

7-CONCLUSION

The statistical analysis presented has the potential to provide new insight into stress concentration factors and into the applicability of statistical models of composites. Attention has so far been restricted to a single statistical model, with F the only parameter describing the stress concentrations. Computational complexities and data limitations make it difficult to estimate more parameters with a reasonable degree of precision. In principle, however, the method may be applied to a far wider class of models.

8-ACKNOWLEDGEMENT

We thank Dr. Michael Pitkethly for providing the data, and Dr. David Clarke for useful conversations on this and related topics. Linda Wolstenholme's work was supported by an SERC research grant.

9-REFERENCES

1. Bader, M.G. and Pitkethly, M.J. (1986), Probabilistic aspects of the strength and modes of failure of hybrid fibre composites. In, Mechanical Characterisation of Fibre Composite Materials (ed. R. Pyrz), Aalborg University, Denmark.

2.Bader, M.G., Pitkethly, M.J. and Smith, R.L. (1987), Probabilistic models for hybrid composites. Proceedings of ICCM-VI and ECCM-2 (ed. F.L. Matthews, N.C.R. Buskell, J.M. Hodgkinson and J. Morton), Elsevier Applied Science, London. Vol. 5, 481-495. 5, 481-495.

3. Harlow, D.G. and Phoenix, S.L. (1978), The chain-of-bnndles probabili ty model for the strength of fibrous materials. I: Analysis and conjectures. II: A numerical study of convergence. J. Composite Materials 12, 195-214 and 314-334.

4. Harlow, D.G. and Phoenix, S.L. (1981), Probability distributions for the strength of composite materials I and II. Int. J. Fracture 17, 347-372 and 601-630.

416

5. Rosen, B.W. (1964), Tensile strength of fibrous composites. AlAA J. 2. 1985-1991.

6. Wolstenholme, L.C. and Smith, R.L. (1989) Statistical inference about stress concentrations in fibre-matrix composites. J. Mat. Sc. (to appear).

7. Zweben, C. (1968), Tensile failure analysis of fibrous composites. AlAA 6, 2325-2331.

8. zweben, C. and Rosen, B. ( 1970), A statistical theory of material strength with application to composite materials. J. Mech. Phys. Solids 18, 189-206.

1.0_ spacing

0.5 _ spacing

TABLE 1

MAXIMUM LIKELIHOOD ESTIMATES OF F

d F 95\ confidence interval

1.0 100 [51,460] 1.5 _ 2.5 120 [75,270] spacing 4.0 170 [97,520]

1.0 80 [57,122] 1.0 _ 2.0 48 [42, 56] spacing 2.5 115 [85,175]

4.0 67 [58, 78]

TABLE 2

JOINT MAXIMUM LIKELIHOOD ESTIMATES OF F AND Xl

d F xd Xl

1.0 80 4.12 4.12 2.0 36 4.08 4.175 2.5 120 3.99 4.114 4.0 46 3.995 4.184 5.0 150 3.88 4.094

2.0 15 4.37 4.47 2.5 16 4.35 4.48 4.0 19 4.26 4.46 5.0 21 4.22 4.45

417

S (mm) , , I , I I ,

A I II 1·5 , I ,

, ,

, , , , , , , , , i , : , , , , : B , , , , , " I , , I I lO

I , I , , I , , I ,

"

( I 0·5 i

0 0·3

Fig. 1 The distribution of tow failures is shown in 4 arrays, each of 7 parallel tows of 1000 carbon filaments in the hybid composite. The distribution is virtually random at the widest spacing (A) but interactions become more intense at the closer spacings, B to D.

n () () (l ()

0

L 1= md J

Fig. 2 Chain-of bundles model

N fi bres

(]

~ 5, 5, 52

Fig 3. Exponential decay of stress ~ in a fibre adjacent to a broken fibre. This is approximated in the model by the step functions shown by the broken lines . The stress level before failure is denoted by ~o.

A STANDARD FOR INTERLAMINAR FRACTURE TESTING OF COMPOSITES

P DAVIES, A. ROULlN"

Ecole Poly technique Federale de Lausanne Laboratoire de Polymeres - 32 Chemin de Bellerive - 1007 LAUSANNE - Switzerland

·Tetra-Pak, SA - Development Department - 1680 ROMONT - Switzerland

ABSTRACT

This paper describes the activities of a task group of the European

Group on Fracture examining interlaminar fracture tests, with a view to

standardization. Despite the increasing use of such tests to characterise

composite material toughness no widely accepted standard test procedures

exist. In order to highlight contentious areas which might hinder the

adoption of a standard, a group of university and industrial laboratories

from 9 countries have been involved in two series of round robin tests on

four different materials, over the last two years. This has enabled test

protocols to be drafted for mode I and mode II tests, which are currently

being examined in a third series of tests.

I - INTRODUCTION

The laminar nature of most polymeric matrix composites gives rise

to a unique failure mode in which cracks propagate through the

interlaminar region. The resistance of a composite structure to delamination

often determines the life of the structure so that a reliable and reproducible

means of measuring delamination resistance is indispensable.

One way of characterising this resistance is by the use of a fracture

mechanics approach. Mode I tests, in which unidirectional composites are

delaminated under tensile loading, have been performed on composites

420

since the late 1960's. A single specimen configuration, the double cantilever beam (DC B), has generally been used, and values of GIc are now commonly

found in material suppliers' data sheets. However, as there are no widely

accepted standards for these tests, the comparison of data from different

sources is hazardous. In addition, there is some evidence of specimen

geometry dependence, particularly in tougher composites, and a

rationalisation is urgently required.

The application of mode II (in-plane shear) and mixed mode tests to

composite materials is more recent and a range of specimens and. techniques

exist.

II - HISTORICAL BACKGROUND

In the USA an ASTM task group was formed in 1981 to study

interlaminar fracture testing and a round robin has been conducted since

1984.

In Europe some attempts have been made to coordinate test

methods, with national initiatives such as the CRAG test standards in the

UK aerospace industry /1/, and a recently published comparison exercise in

France /2/. Wider exercises such as V AMAS also include composite testing,

but at the time of writing no results had been published from this source.

In 1985 the European Group on Fracture set up a task group to look

at standards for fracture mechanics testing of polymers, under the co­

chairmanship of Professor JG Williams of Imperial College, London, and

Professor HH Kausch of the Ecole Poly technique Federale de Lausanne. (This has resulted in the establishment of a protocol for KIc testing of polymers,

which is now being studied by an ASTM group looking to standardize

polymer fracture tests. This protocol has been presented elsewhere /3/.) At

the same time a study of interlaminar fracture was proposed by Professor de

Charentenay, (then at the Universite de Technologie de Compiegne). Two

round robin exercises organised by the authors have been completed since

1986 and results have been discussed at twice-yearly meetings. Eleven

groups participated in the first series of tests and seventeen are involved in

the current (third) round robin. The full list and affiliations of those who

have been involved in meetings and/or testing are presented below.

421

III - PARTICIPANTS

1. EPFL (CH) HH Kausch, P. Davies

2. Ciba-Geigy (CH) M. Fischer

3. Asea Brown Boveri, (CH) D. Neville

4. Tetra-Pak (CH) A. Roulin

5. Imperial College (UK) JG. Williams, AJ. Kinloch, S. Hashemi

6.ICIplc (UK) DR. Moore, RS Prediger

7. Testwell Ltd (UK) N. Trigwell

8. Welding Inst. (UK) G. Hale

9. Univ. Compiegne (F) (FX de Charentenay)

10. Atochem SA (F) B. Echalier

11. UNIREC (F) J. Heritier, S. Thery

12. Rhbne-Poulenc (F) Y. Giraud

13. Aerospatiale (F) D. Lang

14. BASF (D) F. Ramsteiner

15. Fraunhofer Inst. (D) L. Konczol

16. Univ. Hamburg (D) K. Friedrich, H. Wittich

17. Poly. di Milano (I) A. Pavan

18. SINTEF (Nor) B. Melve

19. Warsaw Inst. (Pol) P. Czarnocki

20. Univ. Porto (Port.) A. Torres-Marques, C. Rebelo

21. General Electric (USA) M. Takemori, A. Glessner

22. Texas A & M Univ. (USA) WL. Bradley.

IV - ROUND ROBIN TESTS

The first round robin was conducted on the following materials,

based on a fairly brittle epoxy matrix and a tough thermoplastic. (All tests to

date have been performed on unidirectional material)

(1) Carbon / Epoxy, Vf = 60%, ( supplied by Ciba-Geigy)

Carbon / PES (Polyethersulphone), Vf = 60%, (ICI plc)

422

Groups were simply asked to obtain values of GIc and GIIc in the first

series of tests. A limited number of parameters such as specimen size and

testing speed were given, but other variables were not specified. In mode II

for example some groups used the ENF (Edge Notch Flexure) and others the

ELS (End Loaded Split) specimens.A few of the large number of results

obtained are presented in Figure 1. These tests enabled a number of critical

areas to be identified, such as ;

- discrepancies in data analysis methods,

- propagation mechanisms such as fibre bridging and multiple

cracking, and their implications (initiation v propagation), etc.

In the second round robin participants were asked to follow a more

stringent set of instructions, to allow in particular the assessment of the

feasibility of measuring initiation values of GIc. The materials tested were:

(2) Carbon I Epoxy, Vf = 60%, (Ciba-Geigy)

Glass I PA66 (Polyamide), Vf = 30%, (lCI pIc)

Glass I PU (Polyurethane), Vf = 30%, (leI pIc)

The carboni epoxy was similar to that tested previously but this time

different panels had different starter defect thicknesses. A few results

indicating the influence of the type of defect on initiation values for this

material are given in Figure 2. The glass reinforced specimens were of low

stiffness, so that experience might be acquired with corrections for large

displacements and the use of bonded stiffeners.

As a result of these two series of tests, a test protocol was drafted and

circulated to participants for comments 14/. The evaluation of this protocol

forms the basis of the third series of tests to be performed on two materials:

(3) Carbon / Epoxy, Vf = 60%, (Ciba-Geigy)

Carbon / PEEK, Vf = 60%, (leI pIc).

423

Once again the carbon / epoxy has been chosen as a reference

material, while the carbon / PEEK is the same as that being supplied to the

ASTM task group, with whom we are in contact.

These tests are due to be completed before the next meeting of the

group which is scheduled for May 1989.

v -CONCLUSIONS

In the limited space available a comprehensive analysis of the data

collected is clearly impossible, so only two examples are presented. Within

the task group programme upward of 250 tests have been performed to date,

and these have also led to a considerable number of supplementary tests in

the various laboratories, which have been reported at the meetings. The

protocol which has been drafted reflects the current consensus of opinion

among the participants but the programme is continuing and the evaluation

of this protocol will form the basis for a future publication.

It is in the interests of both material suppliers and composite users

that a single coherent test methodology emerges for interlaminar fracture

testing. It is hoped that liaison with the ASTM task group will encourage the

appearance of such an approach as quickly as possible.

ACKNOWLEDGEMENTS

The authors would like to acknowledge the contributions of all the

groups involved, both in performing the tests and in actively participating at

the meetings.

REFERENCES

1. Curtis PT, (ed.), CRAG test methods, RAE Tech. Report 85099, 1985.

2. Guedra D, Lang D, Rouchon J, Marais C, Sigety P, Proc. ICCM6/ECCM2, 3-

346 London, July 1987.

3. Williams JG, Proc. ECF7, Budapest, Sept. 1988.

4. European Fracture Group, Composites task group, Draft protocol for Mode

I and Mode II testing, 1988. (Available from P. Davies, EPFL, Switzerland).

424

Gic MODE I 3000

• • • • • •

2000 • • • IJ IJ

1000 IJ • IJ IJ

IJ

• • • ~ • • • • ~ • ~ • 0 0 0 0 0

0

Group

Figure 1. Mean GIc values O/m2) obtained from first round robin

(0), (.) Carbon/Epoxy and (0), (_), Carbon/PES, short and long cracks

250

50

r- 1- * c- ,-

~ ,L ~ - jg

!+ 1 -f- ~

~ !~ ~ ~ ~

200

150

100

AL AL PTFE 1x 2x I II 20 40 60

Starter films PrecracK Figure 2. Mean values of Glc (J/m2),at initiation, defined by onset of

non-linearity, for Carbon/Epoxy in second round robin tests.

IMPULSE AND RANDOM TESTS FOR THE MODAL PARAMETERS EVALUATION OF A CFR PANEL

P. GAUDENZI

DIP Aerospaziale Universita di Roma "La Sapienza" - Via Eudossiana 16 - 00149 ROME -Italy

ABSlRACf

This paper deals with the evaluation of modal characteristics of a carbon fiber reinforced laminated panel in the range of frequencies 0-200 Hz. Although composite materials are largely adopted because of their specific properties in structural parts of many kind of systems, their dynamic behaviour is not fully investigated. In this work a single point excitation technique is adopted with impulse and random excitation devices, in order to carry out a modal parameter extraction of the panel under test. Experimental results were processed by a FFT spectral analyzer and finally rearranged to get information on eigenfrequencies and structural damping coefficients. A comparison between the results obtained with impulse and random excitation is also made.

1. INTRODUCTION

Composite materials are frequently adopted in primary structural components. The adoption of composites in aerospace applications has been very extensive due to the high specific mechanical properties performed by advanced composites. For this reason a correct modal characterization of composites becomes a crucial point especially in this field of engineering.

As reminded in the introduction composite structures exhibit a peculiar mechanical behaviour which requires "ad hoc" prediction analytical models. For istance as far as natural frequencies are concerned, a strong shear deformability in the directions normal to the middle plane of a laminated plate influences more and more the values of eigenfrequencies as the order of the mode becomes greater. Such effect, which is relevant also for the first mode if the structure is thick, can be examined from a theoretical point of view with the introduction of three-dimensional analytical models and needs an experimental confirmation. Moreover there are many problems in which dynamic damping is of primary importance. Some significant examples are: a) dynamic response near resonance: in the vicinity of resonance peaks a good knowledge of the damping ratio::; is a dominant factor in the estimation of peak response; b) supercritical aeroelastic vibrations: in aeroelastic phenomena e.g. vibrations of panels at dynamic pressure q greater than a critical value 'k;, a limit cycle is established where the amplitude

426

A is depending on the balance of a cycle average of aerodynamic forces (input) and of dynamic damping, in this case ~ becomes important for design purposes; c) loss of stability in spinning spacecrafts (1). The importance of materials in the damping behaviour of the structure is also well known but there are not sufficient information concerning damping properties of composites. Maybe an ambitious goal could be to control damping properties of a structure by acting on the phisical properties of composites.

2. TIIEORETICAL BACKGROUND

Prediction of the modal parameters of structural systems has assumed in the last years more and more importance. Theoretical and experimental methods and techniques of modal analysis (3) have reached a high degree of complexity in industrial applications too. It is worth noting that, though new advanced procedures are set up, dynamic characterization of structural systems is a diff1cult goal to get. In some crucial cases, such as the damping properties evaluation, experimental tests present very scattered results. For this kind of problems particular care should be payed to test equipment limits, to experimental procedures and related theoretical background.

2.1 The frequency response function

One of the most largely adopted data processing technique is the analysis of frequency response function of the structure under test. As well known, frequency response function H(f), is obtained for the particular output/intput Fourier transform ratio:

H(f)= B(f) / A(f) where f is the frequency. For transient vibration spectral analyzers compute Discrete Fourier Transform (DFf) assuming that the complete transient event is periodic. Alternatively it is possible to derive H(f), both for random and for impulse appoach, from the auto spectra of the response and excitation signal Sbb ' Saa and the cross spectra between the two spectra

Sab and Sba' From those quantities two different estimates of frequency response

function, denoted as HI (f) and H2(f), are available:

HI (f)= Sab (F) / Saa (f) and H2(F)= Sbb (f) / Sba (f) The presence of two equations for the same quantity allows to control quality measurement. To this end coherence function 152 is introduced, defined as follows (2):

152= HI(f) / H2(f) Coherence is a real function, always positive and less than or equal to 1.0; it is a good indicator of the quality measurement, which is well made if coherence is close to unity.

2.2 Excitation techniques

Impulse excitation, generated in this case by an hand-driven impact hammer, consists of a pulse force experienced by the structure during the impact with the hammer. Tips and heads of the impactor were tested in order to obtain a flat frequency content of the pulse in the range of study. The structure under test is relatively small in sizes so we can suppose that in each point all the modes are excited in the frequency range considered. Some difficulties arise for the effects of signal truncation in the time domain. In fact to avoid signal truncation errors, it is necessary to go down to zero with the decaying dynamic response within measurement window. If the mode under concern is

427

lightly damped it is not easy to meet this requirement, being necessary sufficiently long time of signal acquisition.

Random excitation too is characterized by a flat spectral density diagram, though unsatisfactory exciter/structure connection generates irregularities in frequency content. Adopting band-pass filters at the end of generator one can excite only the desidered modes. One of the most evident advantages of this kind of excitation is the easiness and speed of data acquisition from experimental devices.

3. MEASUREMENT PROCEDURES

The experimental tests were carried out on a 300x700x3 mm carbon fiber reiforced panel made by Aeritalia Spa (material standards:BMS 8-212 CLASS 2 TYPE IV STYLE 3K-70-PW). The panel was suspended in a free-free condition by soft springs; the highest rigid motion frequency values were measured resulting at most equal to 10% of the lower elastic ones. The specimen properties were investigated using a single point excitation method, based on the computation of frequency response function. An excitation is produced in one point of the structure; in another point acceleration respose is observed. Measurement chain is made by an exciter (an impact hammer or a shaker coupled with a power amplifier), a charge amplifier connected with one of the two channels of spectral analyzer; the second channel is linked by another charge amplifier with an accelerometer (fig.1).

The basic components of the measurement system used for the tests were: a two channel analyzer SD375 Scientific Atlanta with a broad band and a limited band random generator, an impact hammer PCB 086B01, a power/signal conditioner PCB 480D06, an accelerometer PCB 309A, a shaker BK4809, a power amplifier BK 2712, a charge amplifier BK 2635, a force trasducer BK 8200.

4. EXPERIMENTAL RESULTS

4.1 Tests in 0-200 Hz range

Firstly several tests were performed in the range of 0-200 Hz both for impulse and for random excitation, in order to get preliminary information on eigenfrequencies and to detect modal shapes. In this stage the effects of different kinds of suspension (horizontal and vertical) were also investigated and some simmetry verification tests were done. Concerning the vertical suspension of the panel, accelerometer was firstly positioned in one comer of the panel and excitation was given in all the points of a vertical row of the grid and then in several others points, also to detect structural simmetry. Afterwards the accelerometer was shifted to test the simmetry of the frequency response function matrix. Modal shapes were identified, at least for the first five modes. The first and the other odd modes are of torsional kind, the second and the fourth have bending shape. Results on eigenfrequencies are affected by little scattering both for horizontal and vertical suspension, being frequency resolution in the whole range (0-200 Hz) equal to 0.5 Hz. Referring to coherence diagrams, it is worth noting that low coherence occurs mainly in areas close to antiresonances (fig.3). As far as frequencies and modal shapes are concerned there is good agreement for impulse and random tests as reported in table 1. Apart from the 6th and 7th mode which are coupled and the 1st and the 2nd which appear to be slightly coupled, the others appear to be sufficiently well separated to use the peak amplitude method for modal parameter extraction. This method assumes that all the response is attributed to the single mode under concem, that is the only contribution to response is attributed to the mode whose natural frequency is closest.

428

4.2 Damping factors

After the test performed in the whole frequency range of interest from 0 to 200 Hz, a zoom factor of 10 was introduced. Aim of this stage was the evaluation of the modal damping coefficients as well as of the eigenfrequencies with a greater resolution (0.05 Hz). For the 3rd mode tests were also carried out with a zoom factor of 20. Estimates of modal parameters of the first five modes are reported in table 2 . Damping factor ~ is obtained from real part of transfer funcion according to the well known 'half-power points' method (fig. 2).

Unfortunately large scattering (up to 30 %) affects damping values obtained with different input-output points, both for random and for impulse excitation.

From table 2 it can be noted that torsional modes present higher values of ~ with respect to the bending ones, which are very lightly damped. For these modes low coupling effect can produce large errors in damping ratio estimates.

As far as random excitation is concerned supplementary difficulties arose for damping measurement. In fact attachment devices can produce damping themselves; for the same effect some shifting in measured eigenfrequencies is also possible.

These two aspects can produce an overestimate of damping, which can be extremely dangerous for the consideration done in § 2.

However a fairly good agreement from the two excitation methods is observed, though considering the scattering present in both kind of results. The only significant disagreement is noted for the damping ratio of the 2nd mode, but in this case a low coherence diagram in the region of resonance occurred for random tests.

5. CONCLUSION

A modal parameter characterization of the frequency behaviour of a composite laminated panel was performed adopting two types of excitations. For the structure under test no significant differences emerged from the results obtained following random and impulse approach. Lightly damped modes require elevated zoom factors for correct damping ratio evaluation. Particular care must be payed in tests performed with random approach as far as attachment devices are concerned. Bending and torsional modes exhibit a different damping behaviour. Improvements in modal parameters extraction can be obtained using multi-degrees-of-freedom identification techniques, at present at disposal also as post- processing library software.

6. REFERENCES

1. P.Santini, A. Castellani and A. Nappi - "An introduction to the problem of dynamic structural damping" - AGARD report No. 663.

2. OJ. Ewins - "Modal testing: theory and practice" - Research studies press -Letchworth, Hertfordshire, England - 1984 - pp.130-140.

3. S.R. Ibrahim - "Modal identification techniques assessment and comparison" - Sound and Vibration - August 1985- pp. 10-15.

Mode torsional

30.5 30.5

f (Hz) random excitation f (Hz) impulse excitation

2 bending

36.5 37.0

3 torsional

69.0 68.5

4 bending

101.5 101.0

5 torsional

130.5 131.0

Tablel-Eigenfrequencies (f) for random and impulse excitation, measured in (0-200)Hz.

429

Mode 2 3 4 5

Random Fr (Hz) 30.00 36.12 67.55 101.35 130.06 ~ 10 3 8.42 5.10 5.98 2.46 6.32

Impulse Fr (Hz) 30.26 36.72 67.88 101.17 130.55 ~ 10 3 8.25 2.52 8.83 1.97 7.11

Table 2 - Eigenfrequencies (Fr) and damping coefficients for the examined structure, measured in zoom tests.

SIGNAL GENERATOR

SPECTRAL

ANALYZER

JF LIN 3/A Ave N 10 IE2

V/V

t

o J I

AMPLIFIER

AMPLIFIER

,. )l ]'1

0.0 LIN x HZ 2QO.0 XI lQ7 . 5 HZ ~·2 O. gag iF 1. 91 EI V/ V

Lee B A\'G:: 10

C.O H: <co. 0 X, 197.5 )oC Y(A,) Y (8) 2.6BE-3 V

Fig. 1 Measurement chain.

TF REAL

f

Fig. 2 Half power point method for determining the damping ratio ~

Fig. 3 Coherence function, modulus of frequency response function, output response log-lin spectra (impulse approach, 0-200 Hz)

430

'TTT''!'''! '1' Tllll :ft r '!" hI' T n'l'fl

" -. ... 'tL I'" .. .., .. C :3~oa y""'" TA A C 10 e C 10 II ... " , I I I )-~I I I

( I ) I 1 I \j V I I c

I I I I I T ,WC N I'

'EI F LIN /11.

I I I I I I I I

n I I I I " A.Ve;,. 100 0 .. 0 VIV ti

IJj x. 157.6750 H: YeRE) -l. ZB Y (]).\) 2. B8 II/Y

V I'--j;: an ~E~' ~.,o G 30 1"'19 ... - -~ "" :J 3'" :)"

5 < .. _ W''''' " ".: . " .

~O. J I I I I I I I I I I I I I I I I I

ge.5 LIN X HZ 1 , XI too. 9500 HZ • -079.' DEC TF 3.48 EI V/V

I I I I I I I I I I I I I I I I /T"~ I I I

I I I I I! I I I I I , ,

I I I N/ I I I I I I I I I I I I I I I

;::r"fr fil' r"rll:" ::: 1""1' II'" M' 11111 ';'l~,;;~;;~;~~tf;;;'i~~;;;;t:;:

96.5 LIN X HZ 108.5 mode (!OrslOnal); random approach. X.100.9500l-lZ Y<RE) 8.37 Y(IJol) -3.AOEI V/V

i Fig.5 Coherence function; phase angle, modulus, real and imaginary part of frequency response function for a lightly damped mode (4th mode: bending); impulse approach.

TF UN BIA AVG N 15 5El

I I I I I q . ,

I I I II I ]I. I 111

V/V

IJ\ I 11 I

V \ ~l jJ\ f"-....-o

28.0 LIN X HZ 38.0 X, 33.0000 HZ 1"2 0.962 TF 1.68 V/V

Fig.6 Light coupling effect Ost and 2nd mode).

ABSTRACT

EFFET DES CONDITIONS D'ELABORA TION SUR LE COMPORTEMENT MECANIQUE, STATIQUE ET

DYNAMIQUE DE MATERIAUX COMPOSITES HAUTES PERFORMANCES A MATRICE

THERMOPLASTIQUE SEMICRISTALLINE

c. VEROEAU, A. BUNSELL

Ecole des Mines de Paris Centre des Materiaux - BP 87 - 91003 EVRY cedex - France

Two types of advanced carbon fibre reinforced thermoplastic have been tested and the effects of fabrication conditions on their microstructures revealed. The APC-2 composite is revealed to bond extremely well to the As-4 carbon fibres and this, as well as transverse properties, are improved by heat treatments. This behaviour is shown to be due to the carbon fibres acting as nucleation sites and promoting transcristalline growth. The Ac4o-60 composite showed no such behaviour and subsequently had poor interfacial properties. Heat treatment of the PPS polymer increased crystallinity due to an increase in the proportion of low molecular weight polymer.

INTRODUCTION

A l'heure actuelle, les materiaux composites a matrice thermoplastique connaissent un important developpement

Deux resines ont plus particulierement attire notre attention : - Ie PEEK commercialise par ICI (Imperial Chemical Industries)

composite associe APC2 (Fibres de carbone AS4/PEEK). - Ie PPS commercialise par la PHILLIPS PETROLEUM COMPANY U.S.A.

composite associe Ac4o-60 (Fibres de carbone AS4/PPS).

L'objectif du travail est de relier les performances mecaniques qu'ils soient statiques ou dynamiques aux caracteristiques microstructurales des produits, resultant des divers traitements thermiques entrepris lors de I 'elaboration, en privilegiant l'analyse des phenomenes a l'interface.

432

Dans ce but, nous avons priviligie des modes de sollicitations mecaniques en cisaillement afin de mettre en evidence des modifications eventuelles de la structure du polymere au V01s1nage des fibres, puisqu'il est reconnu que ce type de sollicitation est Ie plus adapte pour evaluer l'adhesion fibre/matrice.

PRESENTATION DES MATERIAUX

Les materiaux composites etudies ont ete elabores sous presse a partir des preimpregnes. Ceux qui ont ete elabores selon les cycles preconises par les fournisseurs sont dits cristallins.

- pour les amorphes, la meme procedure que pour les cristallins a ete sui vie , puis les plaques de materiaux ont ete portees a nouveau a la temperature de mise en oeuvre et pres sees une nouvelle fois sous presse froide de maniere a figer la structure dans l'etat inorganise de la fusion ;

- pour les traites, afin d'eliminer les germes preexistants de la matrice et de rendre Ie role d'agents nucleants des fibres de carbone preponderant, des sejours de longue duree (2h) a la temperature de mise en oeuvre ont ete envisages.

Suite a 1 'elaboration, les taux volumiques Vf (effectues par dissolution chimique du polymere) et les taux de cristallinite X (evalues par DSC ou RX) ont ete determines. Les valeurs critiques sont les suivantes

APC2 Ac40-60 APC2 amorphe cristallin traite

Vf Vf X X X

ETUDE EXPERIMENTALE

61 % ± 1 % 53 % ± 1 % 10 % ± 3 % 34 % ± 3 % 30 % ± 3 %

quelque soit l'etat cristallin " " " "

Ac40-60 amorphe X 5 % ± 3 % cristallin X 51 % ± 3 % traite X 58 % ± 3 %

Une etude enthalpique differentielle effectuee sur un Dsc.4 PERKIN-ELMER a ete suivie d'une etude mecanique statique et dynamique.

Des essais statiques de flexion 3-points avec distances entre appuis variables, des essais de traction transverse et des essais de tractions a ± 45' ont ete effectues, a la temperature ambiante.

Le frottement interieur des materiaux a ete evalue par un pendule de torsion fonctionnant en regime harmonique force a basses frequences.

Analyse enthalpique differentielle

En admettant que l'enthalpie de cristallisation est proportionnelle a la quantite du polymere cristallisee, les taux de transformation sont calcules en integrant les pics de cristallisation. L'evolution des taux de transformation en fonction de la temperature est presentee sur les figures 1 et 2, pour tous les materiaux etudies qui ont ete maintenus a l'etat fondu a des temperatures respectivement

433

300°C et 350°C pour l'AC40-60 et l'APC2 pendant 3 mn seulement.Il a ete adjoint aux materiaux initiaux un echantillon de composite amorphe ayant subi une cristallisation sous contrainte de cisaillement au cours d'un balayage en temperature effectue avec Ie pendule de torsion (cf courbes B). Ces courbes permettent de definir approximativement trois domaines suivant Ie taux de transformation : T

Domaine 1 : 0 < T < 0,3

On remarque que Ie debut de cristallisation pour l'APC2 traite se produit a unetemperature moins elevee que pour Ie composite cristallin alors que dans Ie cas de l'AC40-60, on observe Ie phenomene inverse auquel s'ajoute un tres net ralentissement au debut de cristallisation. De plus, on a une temperature de debut de cristallisation plus elevee pour l'APC2 initialement amorphe cristallise sous contrainte que pour l'APC2 traite suggerantun nombre de germes plus important globalement pour Ie premier composite.

Domaine 2 0,3 < T < 0,7

Dans ce domaine, les courbes sont comparables et ne se differencient que par Ie decalage en temperature provoque dans Ie premier domaine. Le developpement des entites cristallines domine ce domaine et l'evolution du taux de transformation de l'echantillon initialement amorphe et cristallise sous contrainte est tout a fait similaire a celui de l'echantillon traite, suggerant Ie meme type de comportement. Un long seJour a la temperature de fusion du polymere entraine l'elimination des germes preexistants. Ainsi dans Ie composite, Ie role d'agent nucleant des fibres qui a ete montre est favorise par ce type de traitement. On a alors competition entre une cristallisation rapide au voisinage de la fibre et une cristallisation lente dans la matrice. Le surplus de germes cons tate pour l'APC2 amorphe cristallise sous contrainte par rapport au nombre de germes existant pour l'APC2 est donc localise au voisinage de la fibre entrainant une densite de germes a cet endroit tres eleve.

Domaine 3 0,7 < T< 1

Pour l'APC2 cristallin, on observe un net ralentissement du processus de fin de cristallisation lie aux perturbations dans Ie developpement final des entites cristallines induites par la multiplication du nombre de germes et la gene entrainee par la presence des fibres elles-memes.

ETUDE MECANIQUE STATIQUE ET DYNAMIQUE

Etude statique

Des essais de flexion trois points avec distances entre appuis variables ont ete entrepris sur les materiaux composites obtenus dans les etats amorphes cristallins et traites. Ce dernier test simple par sa mise en oeuvre permet grace a une interpretation originale (Hoffman,

434

1985) /5/ des resultats d'avoir acces a des grandeurs physiques telles que la resistance et Ie module de cisaillement ou de traction (G13 - TR et E11 - aR) , la resistance en traction et les energies de rupture. Les resultats obtenus sont rassembles dans Ie tableau 1.

On enregistre pour l'AC40-60 une diminution de la plupart des caracteristiques lors du passage de l'etat amorphe a l'etat cristallin. 11 a ete montre /3/ que Ie passage a l'etat cristallin de ce materiau entraine un etat de contrainte tel au niveau de l'interface qu'on assiste- --lors de la cristallisation au retrait de la matrice qui a pour consequence un affaiblissement de l'interface ; ce que nous avons pu d'ailleurs constater sur les fractographies (cf. figure 3).

L'APC2 presente lui Ie comportement attendu avec une augmentation du taux de cristallinite a savoir une augmentation de la rigidite et de la stabilite thermique. Entre l'etat cristallin et l'etat traite, on observe d'une part une augmentation des resistances au cisaillement et des resistances a la traction transverse mais egalement une diminution de la resistance a la rupture en traction suggerant une amelioration de la qualite de l'interface voire des proprietes de.la zone interfaciale /7/.Les facies de rupture (cf. Figure 4), meme sur celles correspondant a des echantillons d'APC2 amorphe,montrent une pellicule de resine en contact de la fibre suggerant que la qualite d'interface a atteint un niveau tel qu'il est impossible de Ie deteriorer. Les differences de comportement constatees sont donc dues essentiellement a des modifications de la structure du polymere au voisinage de la fibre.

Etude dynamique

Le pendule de torsion utilise est capable de travailler, pour une gamme de temperature allant de -100·C a 600·c, en oscillations entretenues pour de tres basses frequences de 0.5 Hz a 10- 4Hz. Dans la configuration actuelle, la detection est optique et les oscillations sont imposees magnetiquement.

A tres basses frequences, Ie contr6le du spectre de frottement est automatise ainsi que Ie calculateur traitant les donnees. Les eprouvettes ont la forme de parallelepipede rectangle de dimensions : 80 x 4 1 mm3 . Les essais ont ete realises en temperature pour une frequence de 0.1 Hz sur des echantillons de composite APC2 et Ac40-60 obtenus dans l'etat amorphe, cristallin et traite.

Les spectres sont representes sur les Figures 5 et 6; si l'on compare les spectres entre echantillons cristallins et traites, alors que pour l'APC2 on observe un comportement similaire au passage de la transition vitreuse, on cons tate pour l'AC4o-60, l'apparition d'un deuxieme pic ou Ie dedoublement du pic qui correspond a une modification structurale du polymere. Sur Ie spec~re alors que Ie "premier" pics se superpose parfaitement au pic enregistre pour l'AC4o-60 cristallin, Ie deuxieme pic est plus large et decale vers de plus haute temperature comme si une portion de PPS avait atteint un taux de cristallinite encore plus important que celui atteint par Ie

435

PPS cristallin. II a ete montre /2/ qu'avec une augmentation de 1a temperature de traitement a l'etat fondu et du temps de traitement que l'on a une diminution de la masse moleculaire du PPS a haut poids moleculaire.

Cet effet peut entrainer une augmentation du taux de cristallinite ; les chaines etant plus courtes, elles peuvent mieux s'organiser dans l'espace et la cristallisabilite de la resine est donc accrue. D'autre part, il est evident que seule la fraction de polymere au contact de l'02est concernee. Le premier pic se superposant tout a fait au pic enregistre pour l'echantillon cristallin, correspond a 1a relaxation de la phase amorphe de PPS n'ayant pas ete modifiee et Ie second pic a la fraction de PPS degradee pour laquelle on enregistre une diminution de la masse moleculaire. D'autre part, l'augmentation de la temperature de debut de cristallisation corrobore tout a fait cette diminution de masse moleculaire ainsi que Ie ralentissement de debut de cristallisation qui serait lie a une distribution trop importante de masses moleculaires entravant Ie mouvement des chaines.

DISCUSSION

On peut s'interroger sur la presence d'une phase cristalline particuliere au voisinage de la fibre d'autant que 1a presence de transcristallites dans cette region a ete revelee en microscopie optique en transmission (Figure 7) sur des echantillons d'APC2 cristallin.

II a ete montre qu'un refroidissement rapide du polymere (Huson et al. 1984) /6/ ou l'application d'une contrainte de cisaillement (Gray, 1974), /4/ durant la cristallisation de l'etat fondu, ou une forte densite de germes (Burton et al. 1986) /1/ au voisinage de la fibre sont des facteurs qui peuvent favoriser la croissance d'une phase transcristalline. La plupart de ces conditions sont reunies dans Ie cas de l'APC2 cristallin mais surtout dans Ie cas de l'APC2 traite qui a un comportement similaire (du point de vue des cinetiques de cristallisation) a celui des echantillons amorphes cristallisees sous contrainte. Or pour ces derniers, nous avons montre /8/ que l'existence d'un pic aux alentours de 315 0 sur Ie spectre dynamique, temoigne de 1 'existence d'une certaine categorie d'entites crista11ines vraisemblablement de transcristallites. Ainsi la presence de transcristallites au voisinage de la fibre dont l'importance semble etre d'autant plus grande que les echantillons ont subi des traitements a la temperature de fusion semble se confirmer.

CONCLUSIONS

Alors qu'une augmentation de la cristallinite se traduit pour l'APC2 par une augmentation de la rigidite et de la plupart des caracteristiques mecaniques voire par l'existence d'une phase dite "transcristalline" au VOlslnage de la fibre, elle entraine pour l'AC4o-60 un comportement mecanique oppose qui est en fait etroitement lie a la qualite de l'interface de ce materiau.

436

D'autre part si un sejour prolonge a l'air et a la temperature de fusion est tout a fait benefique pour l'APC2 (amelioration des proprietes), il est fortement deconseille sur l'AC40-60 entrainant une degradation du polymere.

En conclusio.1, des polymeres qui semblaient tres proches initialement de par leur structure cristalline a savoir leur reseau cristallin et la conformation de leurs chaines, presentent des comportements tres differents selon les conditions d'elaboration.

BIBLIOGRAPHIE

1 - Burton R.H., Folkes M.J., Mechanical properties of reinforced thermoplastics., Ed. Clegg D.W., Collyer A.A., Elsevier Publish, (1986), pp. 269-294.

2 - Caramaro L., Chabert B., Chauchard J., Influence des fibres de renfort sur les aspects cinetiques et morphologiques de la cristallisation de la mat rice dans un composite PPS/carbone. Note Interne, Universite Claude Bernard Lyon, a Paraitre, (1987).

3 - Davies P., Comportement au delaminage des materiaux composites a matrice thermoplastique". These soutenue a l'Universite Technologique de Compiegne, (1987).

4. - Gray D.G., J. Polym. Sci., Polym. Lett. Ed., 12, p. 645, (1974) .

5 - Hoffman P., These de Docteur Ingenieur soutenue a l'INSA Lyon, (1984) .

6 - Huson M.G., Mc Gill W.J., J. Polym. Sci., Polym. Chem. Ed., 22, p. 3571, (1984.

7 - Norita T., Matsui J., Matsudza H.S., Effect of surface treatment of carbon fiber on mechanical properties of CFRP. Int. Conf. Compo Interfaces, Ed. Ishida H., Koenig J.L., 123, (1986).

8 - Verdeau C., Bunsell A.R., Etude microstructurale de composites hautes performances a matrice thermoplastique. Paru dans Ie recueil des 6emes Journees Nationales sur les Materiaux Composite, JNc6 a Paris, (1988).

Amorphe/Cristallin/Traite Flexion 3 points

G13 (GPa) a R (MPa) TR (MPa)

Traction transverse

a22 (MPa)

E22 (GPa) E %

Traction (± 45)

G12 (GPa)

T12 (MPa)

APC2 Ac40-60

2.3 / 2.5 / 2.5 1.5

1200 / 1800 / 1700 825 / 1000 / 1000

65 / 95 / 105

85 / 60/ 70

8.5 / 11.5 / 10

97

30 / 20/ 15

6.5 / 9.5 / 9

1.66 / 0.45 / 0.7 0.28 / 0.1 / 0.05

4 / 5.5 / 5

125 / 135 / 180

2 / 3.5 / 3

60 / 40 / 50

Tableau 1 Proprietes mecaniques des composites etudies.

l'

z APC 2 0

f- 0,8' < ~ 0:: 0 LL IJl Z 0,5' < 0:: f-

W 0

X 0,4' :J < f-

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0'

330 ; 314 298

l'

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f- 0,8' < ~ 0:: 0 LL IJl z-0.6' < 0:: f-

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Figures

AC40-60

b

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234 218

Courbes donnant en fonction de refroidissement b) Ac4o-60.

437

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439

Figure 5 Amortissement du composite Ac40-60, A une frequence de 0.1 HZ en fonction de la temperature, pour deux etats cristallins : cristallin et traite.

0-1 ("1000) 200

100

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Figure 6 Amortissement du composite APC2, a une frequence de 0.1 HZ en fonction de la temperature, pour deux etats cristallins : cristallin et traite.

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MECHANICAL STRENGTH PROPERTIES FOR ANISOTROPIC COMPOSITES

C.L.D. HUANG

Department of Mechanical Engineering Kansas State University - MANHATTAN, KS 66502 - USA

ABSTRACT

The existence of a strength functional. which is a function of the stress tensor. is assumed. The invariants for the strength function for each class of composites in the third-order approximation are established. Consequently. the strength functions proposed by Tsai and Wu for triclinic and rhombic materials. and by Gol 'denblat and Kopnov for rhombic materials can be obtained readily.

INTRODUCTION

For the purpose of material characterization and design. a rational simple strength criterion for composites is essential and important. As pointed out by Tsai and Wu [lJ. the majority of proposed criteria are limited in their ability to include the correlating stress effects. In order to remove such a limitation. Gol 'denblat and Kopnov [2J proposed a new criterion of strength for anisotropic materials. They investigated explicitly the form of strength criterion for orthotropic materials. In particular, they verified their results for glass-reinforced plastics experi­mentally and showed the suitability of the proposed criterion of strength for practical usage. In this paper, the criterion of strength for anisotropic materials proposed by Gol 'denblat and Kopnov is adopted. We assume the existence of a strength function which is a function of the stress tensor. The invariants for the strength function for each class of composites in the third-order approximation are established. Consequently. the strength functions proposed by Tsai and Wu [1] for triclinic and rhombic

442

materials, and by Gol'denblat and Kopnov [2] for rhombic materials can be obtained readily from the appropriate invariants given in this paper.

1. PRINCIPLE

The general theory of strength functions for anisotropic cyrstals can be established from consideration of the strength function F, where

F=F(o .. )=O (1) lJ

and 0ij is the stress tensor which is, of course, symmetric. The strength function F is required to be invariant under the group of transformations {tij } which characterize the material anisotropy,

F(aij ) = F(oij) (2)

where the transformed stress tensor obeys the following rule

aij = tirtjs °rs .

It is also assumed that the strength function F may be expressed approximately in the following form as proposed by Gol 'denblat and Kopnov [2],

F = (Fijoij)a + (FijktOijOkt)B + (FijktmnOijaktamn)Y - 1

which, in fact, is a third-order approximation of Eq. (1). Th~3) strength function given by (3), which is subjected to the restriction, Eq. (2), for each of the composite classes is determined. In Eq. (3), the Fij , Fijkt , and Fijktmn are strength tensors of rank two, four and six, respectively. The powers a, B and yare material constants. For the case of a = B = Y = 1, the tensors Fij and Fijktmn characterize the Bauschinger effect of materials, and the tensor Fijkt and Fijktmn determine the hyper­surface of the strength function in the six dimensional stress­space.

2. RESULTS

In this paper the classes for composites will be listed by their names together with the transformations defining their symmetric properties. The notation for the transformations is that given by Green and Adkins [3J.

Referring to [4,5], we list the result for each class of composites a set of quantities {I(l) , ... ,I~3)}, each of which is a polynomial of degree three or lower in the six stress components (all = °1' °22 = °2' °33 = °3' °23 = °4' °13 = °5' °12 = (6)' and

443

is invariant under the group of transformations associated with the given class. Hence, the strength functions can be expressed approximately as a polynomial of degree three ·or lower in I,'s with strength coefficients F's. 1

2.1. Triclinic System

Pedial I

Pinacoidal I, C

For crystals having triclinic symmetry, there is no restric­tion on the orientation of the preferred direction; any rectangular coordinate system can be used as a reference frame. Thus, the quanities {Ii} for both classes are

1(,1)., 1 01' O2, 03, 04, aS, 06 .

2.2. Rhombic System (Orthotropic System)

Rhombic-pyramidal

Rhombic-disphenoidal

Rhombic-dipyramidal

2.3. Cubic System

I, ~2' ~3' ~1

!, Ql' Q2' ~3 I, C, ~1' ~2' ~3' Ql'

dO: 1 01' O2, 03

d2) : 1

2 04 , 2 2 05, 06

IP) : 1 040 SO 6

(4)

~2' Q3

(S)

Hexoctahedral (!,~, ~1' ~2' ~3' ~1' ~2' ~3) • (I. !1' !2' !3'

M1, M2) 'V 'V

dO: 1

1(2). 3 '

2 2 + o3(oS + 0 4) +

(6)

444

2.4. Transverse Isotropy

It is supposed that the material is transversely isotropic with respect to an axis x3. Thus, the transformations charac-terizing transverse isotropy are I and B {x' + ix' = e-iu(x + ix2), x3 = x31for all values of u}. u 1 2 1

Therefore, the invariants are

d1): 1 °3' 01 + °2

d2) : 2 2 2 1 °1°2 - °6 , °4 + °5

IP) : 1

{detloi I} (7)

REFERENCES

1 - S. W. Tsai and E. M. Wu, A General Theory of Stren9th for Anisotropic Materials, J. Composite Materials, Vol. 5 (1971), p. 58.

2 - 1. 1. Gol'denblat and V. A. Kopnov, Strength of Glass­Reinforced Plates in the Complex Stress State, Mekhanika Polimerov, Vol. I (1965), p. 70; (English translation) Polymer Mechanics, Vol. 1 (1966), p. 54, Faraday Press.

3 - A. E. Green and J. E. Adkins, Large Elastic Deformations and Non-linear Continuum Mechanics, (1960), pp. 11-13, Clarendon Press.

4 - C. L. D. Huang, The Energy Function for Anisotropic Materials with Couple Stresses - Cubic and Hexagonal Systems, Int. J. Engng. Sci., Vol. 6 (1968), p. 609.

5 - C. L. D. Huang, The Energy Function of Crystal Materials with Couple Stresses, Int. J. Engng. Sci., Vol. 7 (1969), p. 1221.

6 - C. L. D. Huang, Strength coefficients of the cubic polynomal strength criterion for graphite JTA, trans. of the 9th Intl. Conf. on Str. Mech. in Reactor Technology/ Lausanne, 17-21 Aug. (1987), pp. 59-64.

HIGH PERFORMANCE COMPOSITES MADE OF SOLID THERMOPLASTIC POWDER IMPREGNATED FIBER BUNDLES

K. FRIEDRICH, H. WITIlCH, T. GOGEVN, S. FAKIROV'

Technical University Hamburg-Harburg Harburger Schlosstrasse 20 - 2100 HAMBURG 90 - West Germany

'University of Sofia Laboratory for Polymers - bid. A, Ivanov 1 - 1126 SOFIA - Bulgaria

ABSTRACT

Linear flexure-response-studies (EFlex ), through-thickness fracture toughness tests (Kc) and interlaminar mode I and mode II fracture energy measurements (Gico Gllc) were carried out with different laminates of a carbon, aramide and glass fiber/thermoplastic polyamide 12 composite system. Specimens were prepared from fiber bundles interspersed with polymer powder and surrounded by a polymer sheath. The results, which must be considered as prelimary data because of very limited availability of specimen material, reflect an overall good fracture toughness profile of the different laminates tested. This is also highlighted by the fracture surface micrographs achieved with SEM-analysis.

INTRODUCTION

The trend to use thermoplastic polymers as matrices in high performance composites arose from· several disadvantages of the thermosetting resins primarily used by now. Thermoplastics result, for example, in better resistance against interlaminar crack propagation. Additional advantages are: - thermoformability after consolidation, - weldability, - recycling, - new processing techniques, e.g. laser consolidation during fiber placing.

One of the problems with thermoplastics in comparison to thermosets in the past was their high melt viscosity which makes it rather difficult to impregnate fiber bundles in a way that the fibers are wetted well with polymer

446

in the final prepreg or in the laminate. However, several techniques have been developed over the last few years which overcome this difficulty [1]. One method is the guiding of fiber bundles through a fluidized polymer powder vessel. Spreading the individual fibers in the bundle enables an uptake of solid powder particles between the individual fibers. Subsequently, the powder infiltrated bundle is coated with a thin polymer sheath by running it through a die of an extruder. The latter technique was used by Atochem (France, FIT®-material [2]) for producing the intermediate material form needed for the study reported here. With carbon, aramide or glass fibers as reinforcement and polyamide 12 as the thermoplastic matrix it should be investigated how laminates can be made from this intermediate material form (on a laboratory scale) and which fracture mechanical performance can be expected from such composites.

1- EXPERIMENTAL

1.1. Material

The material investigated is a commercial product of Atochem (France), Le. FIT® ( fibre impregnee thermoplastique), consisting out of: - A continuous bundle of 6000 individual carbon (Torayca 6 K, 68 w/o, 0 = 8

Jlm), aramide (Kevlar 49, 47 w/o, 0 =12 Jlm) or glass fibers (Fiberglass 2400 Tex, 70 w/o, 0 = 22 Jlm),

- PA 12 powder between them (particle diameter similar to that of the fibers in order to achieve better infiltration),

- a PA 12 sheath of thickness about 10 Ilm around the whole bundle.

1.2. Specimens manufacturing

Consolidation of this intermediate material form into unidirectional reinforced prepreg sheets was carried out in a hot press using a steel mold (Fig. 1). Consolidation was obtained by heating up the filled mold in the press to 210°C, applying a pressure of 5.5 MPa over a period of 15 min and cooling under pressure down to room temperature within 5 min. In a following step four types of laminates were produced from these prepregs (using the same compression molding technique): (a) [O]n laminates with thickness of about 1.5 mm (only CF/PA) and 3 mm

respectively (Le. 4 or 8 plies) (b) [90,0,90,0]5 laminates (thickness: 3 mm) (c) [0,90,0,90]5 laminates (only AF/PA and GF/PA) (d) [90]e laminates (only CF/PA)

1.3. Mechanical testing

Because of the limited availabitity of test material and of specimen geometry, the different laminates could not be tested comprehensively, i.e. with respect to their total mechanical property profile. Only a few selected tests were carried out at room temperature and moderate loading rates which were

447

expected to provide a good feeling for the potential of this material group in the area of composites with requirement of high fracture toughness.

Laminar flexure response (ultimate stress, modulus) was determined by a three-point flexure test. As the recommended span-to-thickness ratio [3] could not be obtained with the present specimen geometry, a cenain contribution of interlaminar shear deformation affects the resulting flexure data. Fracture toughness (Kc) studies were performed with compact tension (CT) specimens machined from the manufactured composites. The different laminates were used to prepare double cantilever beam (DCB) and end-notched flexure (ENF) specimens for interlaminar mode I and mode" fracture energy tests (Gc) [3].

" - RESULTS AND DISCUSSION

Results of flexure studies on three different laminates out of the CF/PA­system are listed in Table 1. A comparison of the data achieved with the three different laminates at UB = 15 yields a linear trend in the flexure strength values, whereas the flexural modulus shows a progressive trend (Fig. 2). Using unreinforced matrix and fiber data for the prediction of the flexural modulus of the [90,0,90,0]8 laminate, according to the laminate theory concepts [3] gave however, a higher value (Emax = 51.5 GPa) than the one experimentally determined. This is propably due to the span-to-thickness ratio dependence of the modulus at lower UB ratios, as described by Carlsson and Pipes [3]. In the present case, one sample was studied at an UB ratio of about 22 and the measured modulus was clearly higher but the recommended UB ratio of at least 32 [3] was not achieved however. One explanation of the rather low strength may again found in the limited specimen geometry used here. But another reason is definitely the deviation of many fibers from the ideal direction in which they were placed.

From SEM-pictures of the fracture surface of the studied CF/PA specimens obtained by the through-thickness fracture toughness tests, it can be seen that those fibers which were oriented parallel to the main fracture plane look very clean. This reveals rather inadequate wetting of the fibers by the thermoplastiC matrix and by the processing technique used here. The micrographs give funher indication of a high degree of matrix deformation prior to final failure. Kc values were clearly a function of the laminate structure. The individual amounts of energy absorption by the different layers in the laminate can be estimated when convening the Kc values into Gc values (Fig. 3). Each 0° layer absorbs about five times more energy during fracture of the composite than the 90° layers do.

The results of the interlaminar fracture energy tests are listed in Table 2. Under mode I conditions, the results for the CFIPA and AF/PA systems were surprisingly high. For the CF/PA system it is, however, in quite good agreement with the intralaminar fracture energy as calculated from the Kc test of the [90]8 CT samples. Their comparably high values can be attributed to the also very high toughness of the PA 12 matrix itself. In fact, when looking at the fracture surfaces, a remarkable degree of plastic deformation of the matrix

448

between the fibers is visible (Fig. 4a). This is not the case in the GF/PA system because of the lower matrix content and therefore a smaller possibility of plastic matrix deformation between the fibers. In all cases the fibers are barely covered with matrix material, which is a clear indication of the very poor bonding between the components. Also under mode II conditions, the Gllc values of the CF/PA system is in comparison to the GF/PA system higher due to the same reason. Fig. 4b shows the highly deformed matrix under shear loading of the CF/PA system. The effect of the laminate stacking sequence ([O]a or [90,0,90,0]5) is not so significant. But different fiber orientation with respect to the crack growth direction ([0,90,0,90]5) leads to lower values in mode I and more extensive in mode II. This reflects the role of the fibers during fracture, i.e. fiber bridging in mode I and distribution of stress singularities at the crack tip which is only possible when fiber direction and direction of the crack growth is identical.

III - CONCLUSIONS

The continuous fiber bundles, interspersed with fine thermoplastic powder and surrounded by a thermoplastic sheath can be sucessfully consolidated into different laminate forms. The measured mechanical properties reflect the tough nature of the PA 12 matrix used which especially leads to high Kc, Gic and Glic values.

IV - ACKNOWLEDGEMENT

Support of this project by a contract of research cooperation between the German Department for Research and Technology and the Bulgarian Government is gratefully acknowledged (BMFT 227-9211-BUL).

REFERENCES

1 - Cogswell F.N., in Clegg O.w., Collyer A.A. (eds): Mechanical Properties of Reinforced Thermoplastics, Elsevier Appl. Sci., London, 1986.

2 - Atochem Report: Thermoplastic Polymers in Powders for Composites, ATOCHEM, Paris, 1985.

3 - Carlsson L.A., Pipes RB.: Experimental Characterization of Advanced CompOSite Materials, Prentice Hall, Englewood Cliffs, NJ, 1987.

TABLE 1 - Flexural Response Data of CF/PA 12 Laminates

Fiber Fmax Omax EFlex L W B UB orientation {N~ {MPa~ {GPa~ {mm~ {mm~ {mm~ [O]a 475 669 104 40 6.3 2.6 15.4 [O]a 445 627 108 40 6.3 2.6 15.4 [90,0,90,0]5 293 332 19.5 40 6.3 2.9 13.8 [90,0,90,0]5 245 322 20.9 50 6.3 3.0 16.7 [90,0,90,0]5 250 410 46.4 65 6.6 3.0 21.6 [90]a 44 58 4.2 40 6.3 2.7 13.8

TABLE 2 - Interlaminar Fracture Energy Data

Composite system

CF/PA 12 AF/PA 12 AF/PA 12 GFIPA 12 GFIPA12 GFIPA 12

Gk: Fiber orientation (kJ m-2)

Gllc­(kJ m-~

[O]s [O]s

[0,90,0,90)5 [O)s

[0,90 ,0 ,90)5 [90,0,90,0)5

5.0 4.1 3.1 1.4 1.2 1.3

3.1

1.9 1.2 2.0

",--.-,-,..,..,.,.-:-:-:-:,~ Temperature ----Steel Foils

Fiber Bundles

449

Fig. 1 - Geometry of the mold for consolidating fiber/thermoplastic matrix bundles into prepregs and laminates (dimensions: 110 mm x 50 mm).

EF1ex

(GPa)

100

80

60

40

20

(LIB", 15) (Jmax

(MPa) 600

400

0 - _ __ _________ ..L-________ ...1 0

[90]8 [90,O,90 ,O]S {Ola

Fig. 2 - Flexural modulus and maximum flexure stress at break as a function of laminate structure (span-to-thickness ratio '" 15).

450

40~----------------------------'

35

30

25 GC

(kJ / fif) 20

15

10

5

O· · l ayers } Con lr ibullon

- - - - - - - - - - - Co n lribul ion 90 . 1 ~ y«:!S OL---------------~~~~~-

[901s [9 0 ,0 ,90 ,01 5 lO la

Fig . 3 - Fracture energy (Gc) data, calculated from the Kc and flexural modulus data of the [90]s and [90,0,90,0]5 specimens and their extrapolation to a [O]s laminate structure.

Fig. 4 - Comparison of (a) tensile and (b) shear deformation of the matrix, as found on the DCB and ENF specimens, respectively.

EFFECT OF FIBRE VOLUME FRACTION ON TENSILE FATIGUE BEHAVIOUR OF un GLASSIEPOXY COMPOSITE

ABSTRACT

I. PARTRIDGE, P. VIRLOUVET, J. CHUBB, P. CURTIS·

Cranfield Institute of Technology School of Industrial Science, Cranfield Institute of Technology

MK43 OAL CRANFIELD - England "Rae Famborough - Materials and Structures Department

HANTS GU14 6TD - England

The volume fraction of fibres in glass/epoxy unidirectional laminates was varied by inserting layers of the epoxy film (Ciba Geigy 913 resin) into the lay-up. Tensile-tensile fatigue tests were carried out on composites with fibre volume fractions of 69% and 47%. Under a given stress, the composite having the lower fibre volume fraction exhibits significantly longer fatigue lifetime. The micromechanics of damage accumulation and failure has been investigated in order to explain this beneficial effect of resin-rich regions on the fatigue performance of UD laminates.

INTRODUCTION

The work reported here originated from an investigation into the causes of scatter in tensile fatigue data obtained in testing E-glass /epoxy unidirectional (UD) laminates (1). The material used in both studies is of the type used commercially for the manufacture of Westland helicopter rotor blades, in which the lay-up contains a significant proportion of 0° fibres.

rhe initial study concluded that the observed scatter in fatigue lifetimes largely reflects the scatter of static strength in this material. In addition, however, there was an indication of the existence of an 'arrested failure mode' which prolongs the fatigue life under conditions when few fibre fracture sites can be expected, such as at low stress levels or low fibre volume fractions. The present work, therefore, sets out to elucidate the influence of the second of these parameters, the fibre volume

452

fraction, on the micromechanisms of failure and the fatigue lifetimes of the composites.

I - MATERIALS AND METHODS

1.1. Materials

Ciba-Geigy pre-preg Fibredux 913G-E-5-39 was layed-up by hand into 8-ply UD plaques. A second set of 9-ply UD plaques was made in which five plies of the same pre-preg were interleaved with four layers of the 913 epoxy resin film. Both sets of plaques were cured together, in an autoclave, foliowing the manufacturer's recommended cure cycle. The 8-ply plaques were bled on both faces; there was no resin bleed in the (5+4)-ply plaques. The average thickness of the 8-ply pre-preg laminate was 1.15(~ 0.03) mm and the fibre volume fraction, vf ' was determined by the resin burn-off technique to be 69 (t 1)%. The corresponding values for the (5+4)-ply laminate are 0.84(10.07) mm and 47.4 (~2.4)%. The two materials are henceforth referred to by their vf numbers, 69% and 47%.

1.2. Methods

Plain straight sided coupons, 250 x 20 mm, were cut from the plaques using a diamond saw, parallel to the fibre direction. No further polishing of coupon edges was carried out. Following recommended test procedure (2) each coupon was protected by 50 mm long tapered aluminium end tags, which were stuck onto the specimens with Ciba-Geigy Redux 403 adhesive.

Static strength data were obtained from 5 samples using a screw-driven Instron 11954tensile testing machine at a strain

- -i rate of approximately 10 s • An extensometer was used to measure strain.

Tension-tension fatigue tests were conducted on an Schenk 30 Hz resonant machine, using a sinusoidal, constant amplitude tensile stress field with an R-ratio of 0.05. More than ten specimens for each fibre volume fraction were tested to failure (complete separation of the specimen) and a further three of each were tested to estimated half-life. Photographic studies were made of each half-life fatigued and failed specimens. Areas of damage and fracture surfaces were removed and observed on an optical microscope and by scanning electron microscopy using a Stereoscan SP 600 scanning electron microscope (SEM).

Finally, the fibre-resin distributions in both materials were characterised by optical microscopy of metallographically polished cross-sections taken at random in undamaged material.

453

II - RESULTS AND DISCUSSION

2.1. Fibre-resin distribution

The microscopic studies of polished cross-sections clearly reveal the existence of resin rich interlaminar regions in the 47vf material, caused by the resin interleafing. Some fibre migration has, however, taken place and the average thickness of the resin-rich layer is only about 3 fibre diameters.

It must be pointed out that the means of altering the volume fraction by resin interleafing distinguishes the present work from the original study by Barnard et al (1) and those of other authors (3), where different fibre volume fractions are obtained either by the use of different pre-pregs or by altering the amount of bleed during cure. On the other hand, the presence of resin rich layers in laminates is a far from uncommon occurrence and its effect on fibre-bridging and hence on the measured values of the interlaminar fracture energy has already been demonstrated (4).

2.2. Static strength

Results of static tensile tests are summarised in Table I. The failure strengths of the two materials are roughly in the same ratio as the proportion of fibres, as may be expected. The strain to failure is the same for both materials, within experimental error.

2.3. Fatigue results

The fatigue data are summarised in the form of an S-N curve in Figure 1. The beneficial effect of lowering the fibre volume fraction on the fatigue lifetime at any particular stress level is clearly apparent. The split between the two materials becomes even more pronounced if the results are plotted on strain basis.

Our results are in qualitative agreement with those of Puget et al (3), obtained for E-glass epoxy in three-point bend flexural fatigue.

2.4. Damage mechanisms

Model of basic fatigue mechanisms (5) identifies fibre breakage, interfacial debonding and matrix cracking as the microscopic components of cumulative damage development. Our attempts to identify the dominant microscopic features of failure surfaces (6) by SEM and relate them to the principal damage mechanisms have been largely unsuccessful. The problem is one of the complexity of the post-failure appearance in composites.

In glass fibre/epoxy laminates a "macro-photography" technique appears to be more immediately useful (7). The specimen is

454

back-lit and photographs of the observed whitened regions of the specimen and of the intra-ply splitting between fibres are taken at intervals of time.

In simplest terms, the macroscopic differ~nces in the ways in which our two materials undergo damage accumulation and eventual failure can be described by the overwhelming influence of intra-ply 'longitudinal matrix splitting' in the high fibre volume fraction material, in contrast to preferential inter-ply 'delamination' in the low fibre volume fraction material. This delamination dominated failure mode corresponds closely to the 'arrested failure mode' observed by Barnard et al at low stress levels in the initial study (1), when a fibre volume fraction of 45% has been obtained without the use of resin interlayers.

A detailed study of the microscopic causes of stress-induced whitening in glass-epoxy laminates by Nensi (7) provides some pointers to why the presence of resin-rich layers or areas should favour a different mode of damage accumulation to that normally observed in high fibre volume fraction composites, where multiple longitudinal matrix splitting results in a characteristic brush-like appearance of the failed specimen.

Russell (4) found that the resistance to delamination Mode II fracture increased with increasing thickness of interlaminar resin-rich zone in a carbon fibre/epoxy UD laminate. His work also indicated that the same mechanism operates in both quasi-static fracture and fatigue in Mode II. It seems likely that these findings would also apply to the Mode II-dominated delamination failure in our low fibre volume fraction materials. In addition, careful macroscopic observation of the progress of fatigue damage accumulation in our two materials seems to indicate that, at a given stress level, the intrinsic rate of growth of the longitudinal splits is higher· than the rate of delamination propagation along the specimen axis (8). This is consistent with the improvement of fatigue lifetime noted from the 47% vf specimens and provides a qualitative explanation for the findings.

ACKNOWLEDGEMENTS

The authors are grateful to Ciba-Geigy for the gift of materials, to Mr. J.D. Rawles for obtaining preliminary data, to Mr. M. Crook from the College of Aeronautics, CIT, for preparation of the laminates and to Mr. D. Purslow from RAE Farnborough for helpful discussions. This work was carried out with the support of the Procurement Executive, Ministry of Defence.

455

REFERENCES

1 - Barnard P.M, Butler R.J and Curtis P.T, Procs of Third International Conference on Composite Structures, Paisley (I.H. Marshall, ed.), (1985) 69-82, Elsevier ASP

2 - CRAG Test methods for the measurement of the engineering properties of fibre reinforced plastics (P.T. Curtis, ed.), RAE Technical Report 88012 (February 1988)

3 - Puget P., Fiore L., and Vincent L., Comptes Rendus de JNC5, Paris, (9-11 septembre 1986) 715-728, PLURALIS

4 - Russell, A.J., Polymer Composites, 8, No 5 (1987) 342-351

5 - Talreja R., J. of Composites Technology and Research, 7, No. (1985) 25-29

6 - Purslow D., Composites, 17 (October 1986) 289-303

7 - Nensi T., MoD Final Report, D/ERIa/9/4/2064 071 XR/MAT (1988)

8 - Virlouvet P., MSc Thesis, Cranfield Institute of Technology (UK), 1987

69 47

Table I - Static tensile strengths of un glass/ epoxy laminates with two different fibre volume fractions, vf

Stress to failure (MPa)

1337 :!: 63 1039 :!: 63

Strain to failure (%)

3.0 ± 0.3 2.8 t 0.1

Young's modulus (GPa)

46.5 ± 0.3 38.3t1.3

456

-' 0 ..

-' Zo c "'" 3 c:r 111 -,

e -,..., -,..., -111 V! -e -AI -C -, -' 111 <:> ..

-0 ...

Maximum t!nsile stress (MPa) -' '" I.ro.I ~ U1 0- --.l 0 0 0 0 0 0 <:>

0 <:> 0 <:> <:> <:> <:> <:> I I I I I I I

~

0 0 ~

0 • r-~ • l- • r-~

0 • I- 0 • r • 0

0

r • • • • I-

0

l- • l-

0 r • t-l- • --

• 0

- 0

• - 0 •

0 < <

- ~ --• - • 0- ~ - ..0 --.l

~ ~~ l- • r-

Fig. 1 - S-n curves obtained in tension-tension fatigue (R=0.05), at 30 Hz

oc <:> 0

ABSTRACT

MATRIX SELECTION FOR GRP FATIGUE LOADED STRUCTURES

A. GUEMES, J.A. GLEZ-VECINO, M.A. CASTRILLO"

ETSI Aeronauticos Ciudad Universitaria - 28040 MADRID - Spain

"ETSt'lndustriales Crta. Castiello - GIJON - Spain

It IS evaluated the static and fatigue properties of fibre dominated glass reinforced plastics with polyester, vlnylester and epoxy resins, In plain specimens and in presence of holes and bolts. In spite of the qUite different aspect of fracture surfaces, matrix qual ity do not Influence the static nor dynamic strength. In bolted jOints, failure mode shift from bearing fal lure In static tests to tenSile failure under dynamiC conditions, meaning that bearing strength does not degrade at all, or at least do It at a much minor rate that tensl Ie strength. Fatigue degradation rate IS more severe for plain that for notched specimens.

I NTRODUCT I ON

Wind turbine blades, or spring leafs for cars, are good examples where a low cost, high quality composite IS needed. Continuous glass fibre compOSites, With appropriate orientation and high fibre content, can afford strength higher than 600 MPa, making this materials good candidates for many engineering appl icat,ons. Some uncertainties arise from subjects I iKe durabi 1 ity, behaViour in presence of notches and bolted JOints, etc. High performances matrix other than conventional polyester resins, are sometimes selected, In the feeling that they could give a higher confidence.

OUr actual understanding on the fatigue degradation processes In multidirectional laminates /1/, or on the effect of holes and notches /2/, is that fai lure initiates by microcraKs and delaminations at the matrix, which Will suggest that matrix qual ity will dominate or at least Influence these properties of the compOSite. This could be true for CFRP, but it can not be sustained by experimental eVidence on GRP; matrix properties are not reflected on the composite properties; If any, there IS only a slight advantage for the cheaper polyester

458

resin, In comparison with higher qual ity resins. After Its own experimental results and literature review, Mandell

/3/ concludes that matrix properties have not a significant contribution to tensile fatigue performances of GRP plain specimens, showing for a variety of materials a degradation rate of approximately 10 % of their Initial ultimate tensile strength per decade of fatigue cycles, with same exceptions occurring for fabric reinforcements.

Ref. /4/ shows that compressive fatigue behaviour of GRP is at least as good as tens I Ie behaviour.

Referring mechanically fastened compOSite jOints, a recent review /5/ shows that a large number of parameter has to be considered, that, at the present time, analytical approach has to be supported by tests, and that experience on a fibre/matrix system can not be directly traslated to other systems. Design approach for static conditions can be found at /6/.

This paper reflect data obtained on a variety of GRP systems, in a fi·bre dominated (60% O' ,20Y. 90' ,20Y. !45·) configuration, under static or cycl ing tensile loading, in plain speCimens, and in presence of holes (by-pass loads) and bolts (bearing loads). OUr results confirms that matrix has a lesser influence that is supposed to do.

I - EXPERIMENTAL METHODS

Materials used were E-glass fabrics, plain weave (269 gr/rrf) and unidirectional (420 gr/~). Resin systems were a conventional isoftal ic polyester (Cronolite 1112), with benzoyl peroxide catalyst (0.5 Y. by weight); a vinylester resin (Derakane 411-45) with the same amount of catalyst and a tough epoxy resin (Bepox L265). By hand lay up, laminates in the configuration (0, (!45) , (O,90ns. where inner parenthesis means for crossed cloths, where prepared and cured in a hot plate press under a pressure of 3 KPa and 90· C of temperature. Fibre volume, measured by calcination and weighting, was qUite uniform 6O!2 Y. in all the cases. Porosity was negligible < 1 Y.. After machining the specimens at the size given at fig. I, with aluminium tabs, static and dynamic tests were run on a MTS 820, under load control at room temperature conditions. Frequency selected was 10 Hz and only a sl ight heating «10' C) was observed at same cases. Five identical specimens were tested in each condition.

In mechanical joints tests, net fit and washers (11 mm) were used, with a high clamping torque (; 6 Nm). Hole diameter to end distance and speCimen width was maintained as 0.25, promoting bearing failure at static conditions After cycl ing loading, failure mode shift to a tensile mode, starting at the upper edge of the washers and not at the minimum cross section; GRP material inside the washers was intact at failure. After fatigue failure of one side of the speCimen, its other side was tested for residual strength without dismounting. it gives sl ightly higher values than previous statics tests, and failure initiates outside the washer, starting at the contours and propagating in a mixed tensile-shear mode; a quite different morphology than purely static test.

It was evident that the high pressure of the washers minimize the deformations and, consequently, the degradation under cyclic loading of the compoSite inside its circle. Consequently, test of by-pass loading were done with the free hole and with a 'idl ing bolt'; the bolt does

459

not transfer any load, but sustain a tllgh pressure on the washers. Its beneficial effect under cycling conditions are eVident, with a fatigue I ife fourth fold higher than free hole.

I I - RESULTS AND DISCUSSION

Table shows the results of the static tests, and its data reduction to a Welbul I curve (f(x)=exp(-(x/S)~a). In· short, advantage of any Kind of resin IS not found, the configuration is almost notch insensitive (SCF<1.2), and constraining the edge of'the hole has none effect; for the width to d 11ITleter rat lose I ected, In bo I ted jO i nts local bearing Will be by the fal lure mode.

Spreading of data, IS larger for the epoxy resin, which could be interpreted as If, due to the better fibre-matrix adheSion, minor defects can not be rounded off and propagate catastrophically. Fracture morphology differs considerably, from the brush I IKe aspect of the polyester composite to Irregular cracK In the epoxy laminate.

Table I I and I I I shows the life spans under a cycl ing loading of 500/5000 Nand 800/8000 N, the S1lTle for al I laminates; maximum stresses are, approximately, 0.25 and 0.4 U.T.S., respectively. (To remark that Width of plain specimen IS sma I ler than holed specimens, so net stresses are the S1lTle In both cases.)

Fig. 2 shows some surv I val probabll It les. The most fact is the beneficial effect of constraining deformations as expected, I tiS found that degradat I on r'ate I S lesser

significant In the hole; for notched

specimens than for plain speCimens, as shown in fig. 3, mainly with idl ing bolts.

TrYing to adapt the average values to the Mandell's formula, S/UTS=l-(lg N)/b, gives 'b' values between 5.66 for plain specimens to 10.3 in Idling bolts. Agreement With previous publications /3/ IS remarkable, but subtantlatlon In the low stress regime IS needed.

Static and cycliC transfer loading by bolted JOints were also tested. Mean bearing static strength was 574 MPa (casually coincident With urS) and, after fatigue, reSidual strength grows up, In some cases more than 20 I.. As mentioned, there was a transition In failure mode. For gross stres levels In the laminate of 100 MPa, average life was 137000 cycles; for 50 MPa, none of five specimens fal led after two millions cycles.

I I I - CCNCLUS IONS

Aga In, I tiS ver I f,ed that GRP degrades monoton I ca I I y under cycling loadp, and that In presence of holes, fal lure has the characteristic of 'sudden death', meaning that reSidual strength IS not affected and can not be taKen as an Indicator of fatigue damage. Resin qual ity does not have any substantial ,nfluence on thiS rate.

Constraining the edges of the hole does not modify the static behaViour, excluding the obVIOUS Increase In bearing strength /7/, but

Influence substantially the I ife time. In bolted JOints, the fal lure mode WI I I be different ,n static

that in dynamiC conditions.

460

RESIN Polyester Vlnylester Epoxy

SPECIMEN P H I P H I P H I

STATIC 580.5 516.7 529.5 597.8 563.3 510.3 580.5 523.1 500.8 580.5 523.1 545.4 637.9 572.2 545.4 606.0 531.4 535.9

STRENGTH 588.2 545.4 558.2 650.7 576.7 555.0 613.7 539.0 539.0 614.3 548.6 559.8 653.9 581.8 561.4 616.9 556.6 542.2

(MPa) 646.2 564.6 561.4 666.6 593.9 606.0 620.7 564.6 545.4

a 22.9 34.9 67.3 38.9 58.6 18.8 65.7 38.6 57.l 13 594.2 535.9 549.6 638.6 574.8 547.3 606.4 539.4 531.7

TABLE I Static strength. Experimental data and Weibull parameters.

RESIN Polyester Vinyl ester Epoxy

SPECIMEN P H I P H I P H I

FATI~..JE LIFE 1433 1343 1578 1225 1086 1466 1677 1403 3755 (cycles) 1903 1354 1922 1248 1093 1626 2202 1569 3878

0 =.40 llTS 2518 1378 2000 1538 1160 2670 2716 1661 4640 max 2814 2162 2206 1609 1634 2741 2726 1750 4792

R = 0.1 3189 2261 2598 1642 1668 3808 2757 3195 6786

a 4.4 4.4 6.6 10.2 5.5 3.2 7.9 3.0 4.4 13 2196 1571 1964 1419 1252 2166 2344 1660 4382

TABLE I I Fatigue life at cyel IC stress 0.4 U. T. S.

RESIN Polyester Vlnylester Epoxy

SPECIMEN P H I P H I P H I

FATIGUE LIFE 12210 8892 52846 8339 4560 26869 4874 7825 35792 (cycles) 13893 13443 62820 8532 5930 44343 9498 8786 36115

0 =.25 llTS 16099 13531 62954 8683 6002 44725 10742 19906 47181 max 19555 16263 69082 8842 6022 52583 10798 22610 67831

R :: 0.1 21436 16740 76284 9660 6614 54803 16859 29815 71202

a 5.5 6.3 9.4 17.9 11.8 6.0 3.0 2.3 3.8 13 15686 13171 62881 8663 5717 42586 9163 14450 46840

TABLE I I I Fatigue life at eyel Ie stress 0.25 U. T. S.

P : Plain specimen; H : Hole specimen; 1 : Idling bolt For H and I, stresses referred to minimum eross section (net stress).

~I 1-- -3+- - '~ - - -11-- -1 . ~±!l

{! -- I~ ---~6- t/-- -- J

~ I t-- m ~ -,4: IJ --I r 60 - 1

F - ,- ,- 0 -,- ' " .. _-_ .- -·0 -I 110 .--~

Fig. 1

Specimens dimensions

-'. -'. ~.... pl h ZZ;:~:~ .. , I ..... :;. '.., .... ~~ ", :-. \ , )--. "., up, tt.. . ..,

. \ . ep p,-, ' """ .... .... '. U'(p, h_\, ~,-PJP _ \\, " , '. ... ·. up p

. \ '( 'I. t \ \, .. , 1\ \. ' \ ~. , \ep, 1 \. . . '. \

\\, up , 1\:'., \ '/ ' ' . ' '\ " / "'\ ~" \' '. , "

-~----, ..... ......

461

plain

notched

idling bolt

load t ransfe r

·'. ep i \uo, i

'\.~ 1\'-I

\ I '. I,

.- \. - I \" \ I I

,\Y _ " \_ \ ~ _I.. \,., ... , -t''. \.-l II'" \ "\ \.'1 "\

r., I\,\ \ . . \ " "" ~ I', I \ 1 1 ,\', I,'" I'l l \ _ '" _". 1,_

--!. ', ..L..\ , - , - -',', I \ I I "'. , '1"1 •• ',

iJ 14~ \ . 1 " \ I, "I _ \ .

[ _ _ \\, " 'VI ()max - If \ I, \', " .... I, ',I. 0 . 4 UTS I ' I- ! ',., J '\ ..

-\ ~ ~- - - '. '\~. , . , \ . \\ \, \1 \ ~ \ , ',\

\, '\ ,:".ujP,p I' \ I, '., \ \ " .. '\ \ " ' ...... \ . \~~ ...... . '" ... \ . ", .... .. . .

a-max = \l '1 -! 0 . 25 UTS . I - - - "~ I

1'\ I, .. \ \. \

1000 5000

Cycles

Fig. 2

Survival probabilities of polyester and epoxy laminates.

462

18200 N

5000 N

·--·-4-····_·· .'. f p 8850 N

'---.,

8000 N

5000 N

Fig. 3

ACKNOWLEDGEMENTS

log N

"'1 h P

f.poxy »to r-i n

_ .. ...

S-N

i ; P h

·f i

., i

The work reported in this paper was supported under a contract for development of wind turbine blades by the 'Instltuto de EnergTas Renovables', Madrid. The materials were suppl led by CrlstalerTa Espanola, Dow Chemical and Galresa. The authors wish to express their appreciation to each.

REFERENCES

1 - TaireJa R., Fatigue of Composite Materials, Tecnomic Publ. Co ( 1987)

2 - Awerbuch J., in Fracture Meehan I cs Maps for Compos I te Materials, ESA - OR (P) - 2017, Yol. 4 (1985)

3 - Mandel I J. F., Huang D. D., McGarry F. J . . Composites Tech. Review, YOI. 3, N· 3 (1981) 96 - 102

4 - Conners J. D., Mandell J. F., McGarry F. J., 34th Annual Conf. SPI, New Orleans (1979)

5 - Poon C., In BehaViour and AnalYSIS of Mechanically Fastened Joints in CompoSite Structures, Madrid, AGARD - CP - 427 (1987) 1 - 28

6 - Matthews F. L., In Composite DeSign, ThinK Composites Ed. (1987) Chapter 18

7 - Smith P. A., AShby M. F., Pascoe J., Model ling Clamp-Up Effects In Composite Bolted JOints, Journal of Composite Materials, Yolo 21 (1987) 878 - 897

FISSURATION CRACKS

Chairman: Dr D. C. PHilLIPS A. E. R. E. Harwell

INFLUENCE OF THE FIBRE-MATRIX INTERFACE ON THE MATRIX CRACK DEVELOPMENT IN

CARBON-EPOXY CROSS-PLY LAMINATES

J. IVENS, M. WEVERS, I. VERPOEST, P. de MEESTER

Department of Metallurgy & Materials Engineering Katholieke Universiteit Leuven, De Croylaan 2 - 3030 LEUVEN - Belgium

ABSTRACT

Composite materials need a good stress transfer between the fibres and the resin. Therefore, carbon fibres are surface treated. The influence of this surface treatment on the mechanical properties and the damage development in a cross ply (OZO,90 Z0)S carbon epoxy laminate during monotonic tensile testing was investigated. It is shown that the mechanical properties of a cross ply laminate improve for a low surface treatment level, while they decrease strongly for higher treatment levels.

I-INTRODUCTION

In todays industry, composite materials are becoming more and more important. This great interest in composite materials has resulted in research for new resins and new fibres. However, this has lead to the fact that the interface has become very important in the mechanical properties of the composite/I,Z/.

In order to obtain a good load transfer between fibre and resin, good adhesion between fibre and matrix is necessary. According to Goan and Prosen/3/ the two predominant parameters for the adhesion between a fibre and the resin are:

- the surface roughness of the fibre; - the surface tension of the two materials;

Although the surface roughness also contributes to the adhesion, the surface tension is the most important factor. It is influenced by the presence of chemically active groups on the surface of the fibre.

466

Many authors/4,5/ have proved that an oxidative surface treatment increases the amount of chemically active groups on the fibre surface.

When the mechanical properties of these surface treated composites were tested, many authors focussed on ILSS testing. Only a few authors/6/ compared different surface treatment levels during real mechanical testing like monotonic tensile testing and impact. On the other hand, very little is known about the influence of the surface treatment level of the carbon fibres on the damage development in a real laminate.

The goal of the research done at the Department of Metallurgy and Materials Engineering of the K.U.Leuven is to get a better understanding of the influence of the interface on the mechanical properties and the damage development in carbon fibre reinforced composites.

2-EXPERIMENTAL PROCEDURE

2.1.materials

All tests were carried out on crossply laminates with stacking sequence (02,902>S'

The resin used for the tests was an epoxy resin HG9101 of Hysol Grafil Co. The mechanical properties are:

Tensile strength: 40 MFa Tensile modulus 3 GPa Failure strain 1.3 %

The fibre was an mechanical properties:

Tensile strength: Tensile modulus Failure strain

XA fibre of Hysol Grafil with the following

3000-3300 MFa 220- 240 GPa 1.2-1.4 %

The fibres received an oxidative surface treatment. The duration of the treatment was varied in order to obtain different surface conditions:

100 %: the standard surface treatment of the XA fibres. o %: untreated fibres.

10 %: treatment time 10 % of the standard treatment. 400 %: treatment time four times the normal treatment time.

The surface treatment level has a great influence on the chemical activity of the fibre surface/7/. This causes an increase of the surface energy, as proved by Robinson et al./8/. Due to a higher surface energy, the wetting of the fibre is improved and the active groups can form bonds with the resin.

2.2.Cure cycle

The (002,90 02>S cross ply laminate was cured in a press at 175°C during 1 hour, after a heat-up period of 1 hour, with a pressure of 590

467

kPa. The overall fibre volume fraction was measured. It was 61 ± 1% for all specimens.

2.3.The testing procedure

The cured plate was cut into specimens which were tested on' a tensile testing machine Instron 1196 (250 kN). The strain was measured with strain gauges.

The damage development was investigated with the edge replication technique /9/. In order to investigate crack nucleation and growth, the specimen edges are sandpapered up to 1200 IJ.III and then polished on a Petrodisk-M. To produce a replica of a polished edge, a piece of cellulose acetate tape is immersed and softened in acetone. The tape is subsequently pressed onto the specimen edge and allowed to harden. After removal the replica gives an exact copy of the edge of the specimen.

3-RESULTS AND DISCUSSION

The mechanical properties, measured during tensile testing, are presented in table 1.

3.1.Tensile Stiffness

From figure 1 it can be seen that the tensile modulus improves with increasing interface strength up to 60 GPa. Then it reaches a plateau somewhere between 10 and 100 %.

If we consider an ideal laminate, we can use the rule of mixtures to calculate the stiffness:

Eoo - Vf·Ef + Vm.Em - ±130 GPa E900 = (Vf .Ef - 1 + Vm.Em- 1)-1 = ±6 GPa

(0 0 layer) (90 0 layer)

The transverse modulus of the fibres Ef ' is assumed to be 10 GPa. Using laminate plate theory, the modulus of the crossply laminate can be calculated: ECp· 68 GPa. From the results in figure 1, one can clearly see that this value is not reached. This is caused by two reasons.

First, in the 90 0 layer, the fibres can contribute to the stiffness if the bonding with the resin is adequate. The bonding is of course influenced by the fibre surface treatment. If we assume that the untreated fibres are not bonded to the matrix at all, then the stiffness of the 90 0 layer becomes 1 GPa, because the fibres do not contribute to the stiffness. We can assume that the 90 0 stiffness increases up to 6 GPa for the standard treated fibres.

Second, the 00 fibres can only contribute if the load is transferred to the fibres. This load transfer is mainly done by shear load introduction. Stress transfer requires a good interface. Therefore, fibres which are not bonded with the resin will not be loaded. Gresczuk /11/ incorporated this in the rule of mixtures by using Vfb ' the fraction of bonded fibres. Based on the assumption that,

468

if all fibres are bonded, the rule of mixtures is exact, one can calculate the fraction of unbounded fibres. This is shown in figure 2 for the 0° layers.

3.2.Failure stress and strain

In figures 3 and 4 the influence of the treatment level on the maximum stress and strain are presented. Both figures have the same form: they reach a maximum at about 10% and decrease towards the value of the untreated specimens for 400% of the standard treatment time.

The increase is of course caused by the better bonding: more fibres will be stressed, because Vfu is lower (figure 2). However, due to a higher surface treatment, tbe composite behaves more brittle. Growing cracks are not blunted by debonding and therefore remain very sharp. The stress concentration at the crack tip is thus very high, and the crack will grow very fast, even causing fibre splitting. At the 0°_ 90° interface, no blunting will occur, and therefore the 90° cracks will initiate fibre failure in the 0° direction. This leads to a decrease of failure stress.

This can also be seen in figure 5, which presents the results of the edge replication technique. The crack length is the largest for the specimens with untreated fibres, while the specimens with the standard treated fibres contain the lowest amount of cracks. In the specimens with the heavily treated fibres, there are no cracks until a strain level over 0.2 % is reached. However, from that point on, the cracks nucleate and grow much faster than ip the other samples.

For the specimens with a low treatment level, the bonding is not too good. Therefore, the material contains a lot of flaws, which lead to a high amount of cracks. In the specimens with highly treated fibres, the amount of flaws is very low. Therefore, the damage initiation occurs at much higher stress levels. However, once a crack is nucleated, it can grow very fast, as explained earlier. The stress concentrations at the crack tip can lead to fibre splitting.

4-CONCLUSION

From the results it can be seen that the fibre surface treatment does have an important influence on mechanical properties. In order to reach the properties which can be calculated from the rule of mixtures, some surface treatment is needed.

However, one must be careful with the surface treatment level, because at high treatment levels, the composite behaves very brittle: cracks nucleate at a higher strain level, but their growth rate is much higher. This leads to a sudden brittle failure at a low strain level. Therefore an optimum treatment level is obtained for rather low surface treatment levels.

5-REFERENCES

1. S. Lehmann et al.,SAMPE Quarterly, (April 1985) 7-13.

469

2. I. Verpoest et al. in "Academia Analecta", (1989~

3. J.C. Goan and S.P. Prosen in "Interfaces in Composites", (ASTM STP 452) (1969) 3-26.

4. P. Ehrburger and J.B. Donnet in "Handbook of Composites vol. 1, Strong Fibers (Watt, Perov,Ed.) (1985) 577-603.

5. H. J!ger in "Damage Development and Failure Processes in Composite Materials" (I.Verpoest, M.Wevers, ed.) (1987) 79-89.

6. T. Norita et al. in "Composite Interfaces" ,(Ishida, Koenig,ed.) (1986) 123-132.

7. R. Robinson et al. in "High Tech - The way wto the nineties", (1986) 299-310.

8. R. Robinson et al., research paper 9. M. Weyers in "Identification of Fatigue Failufe Modes in Carbon

Fibre Reinforced Composites, part 1 and 2",(K.U.~euven) (1987) 11. L.B. Gresczuk in "Interfaces in composites", (ASTM STP 452) (1969)

42-58.

6-FIGURES AND TABLES

E !Tm Em v (GPa) (MPa) (% )

o % 49 653 1. 07 0.053

10 % 56 805 1.41 0.035

11100 % 63 775 1.24 0.04 II 11400 % 61 580 0.96 0.069

Table 1. Mechanical properties of the (002,90 02)S carbon-epoxy laminate in a tensile test. (E: tensile modulus; !Tm: tensile strength: Em: failure strain: v : Poisson's coefficient)

70

Ii' Il..

~ ~ 60

rJ) rJ) UJ Z SO U. u. i= rJ)

100 200 300 400 500 SURFACETREATMENT{%)

Figure 1. Influence of the fibre surface treatment level on the tensile modulus.

470

0.3.------------------,

:0 r'l 2. 0.1 I ~'_----------->

0.0 +--~__,r__~-__r-~-_._-~-.,._--I o 100 200 300 400

SURFACE TREATMENT (%)

Figure 2. Influence of the surface treatment level on the fraction of unbound fibres in the 0° direction.

..... 900 aI ~

== ...... 800 en en

UJ a: 700 ... en UJ a: 600 :) .J C( SOO U. 0 100 200 300 400 500

SURFACE TREATMENT (%)

Figure 3, influence of the fibre surface treatment on the failure stress.

1.S

~ 1.4 ...... Z C( 1.3 a: ... 1.2 en UJ 1.1 a: :) .J C(

1.0

U.

100 200 300 400 500 SURFACE TREATMENT (%)

Figure 4, influence of the fibre surface treatment on the failure strain.

471

-- 2 :::E :::E --0-- 0% -:::E 10%

:::E -- - 100%

::z:: ---+-- 400% ~

" Z W ...I

~ 0 ~ 0 a: 0 0.0 0.2 0.4 0.6 0.8 1.0

STRAIN (%)

Figure 5: results of the edge replication technique: total crack length per mm specimen length as a function of the strain for the four different surface treatment levels.

ANALYSIS OF THICK LAMINATES USING EFFECTIVE MODULI

C.T. SUN, W.C. LlAO

School of Aeronautics and Astronautics Purdue University - 47907 WEST LAFA YETTE -Indiana - USA

ABSTRACT Three-dimensional effective moduli are used in the analysis of thick laminates

consisting of a large number of identical sublaminates. A global-local method is developed in which the effective modulus theory is used to model the thick laminate except for the local region where accurate stresses are desired. Evaluative examples involving free edge stresses and interlaminar cracks are presented.

INTRODUCTION When three-dimensional stresses are present in a laminate composite, classical

laminate theories (including high order plate theory) are not adequate, and three­dimensional stress analysis must be performed. If the laminate consists of a large number of laminas, it is impractical to consider each individual lamina in the analysis.

In actual applications, many thick laminates possess a certain periodic stacking sequence in order to avoid warping due to presence of curing stresses. If the characteristic length of deformation of the global laminate is large compared with the periodicity, then the nonhomogeneous properties over each typical sublaminate may be smeared out and effective properties used. Thus, the whole laminate can be effectively represented as a homogeneous anisotropic solid. This approach was' investigated by some authors for layered media [1-4].

The effective modulus theories proposed in [1-4] cannot be used where high stress gradients are present; e.g., stresses near free edges, and small cracks in a laminate. A global-local approach has been employed by several authors to deal with such a situation [5,6]. In general, this approach involves modeling the local volume of interest (where high stress gradients exist) by using the exact constitutive relations, and modeling the remote volume by using the effective moduli.

In this paper, the global-local approach in conjunction with the effective modulus theory developed by Sun and Li [3] is employed to analyze the free edge

474

stresses and interlaminar cracks in thick laminates. Full scale finite element analysis results are obtained for evaluative purposes.

I - EFFECTIVE MODULUS THEORY Consider a thick laminate which is formed by repeating sublaminates. Each

sublaminate contains N orthotropic fiber composite laminas of arbitrary orientations. The essence of the effective modulus theory developed by [3] is the long wave assumption regarding the basic deformation characteristics in the sublaminate, i.e.,

- k- k- . .Ie Exx = Exx ,Eyy = Eyy , Yxy =Yxy

-k - k - k azz = azz ,axz = axz ,ayz = ayz (1)

where (1ij ilIld Eij are the effective (average) stress and strain, respectively, and superscript k indicates the kth lamina in the typical sublaminate. The effective stress-strain relations are given by

(2) where

- ------ T - ------ T { a } = { axx ayy azz ayz azx axy } , {E} = { Exx Eyy Ezz Yyz Yzx Yxy }

The explicit expressions of the effective elastic constant Cij'S can be found in [3].

If the laminate contains a single composite system and has a balanced layup, then the effective elastic constants reduce to

N Cij = L VkC~) i,j = 1,2,3,6 (3)

k=1

[k=N ] cpq= LVkC~/L1k /L1 p,q=4,5 k=l

(4)

L1 = [f VkC~/ L1k] [f vkcgcl L1k]- [f vkc2'~/ L1k]2 k=1 k=1 k=1

(5)

L1k =c~ c~~ - [c2'~ r (6)

where Vk is the volume ratio of the kth lamina in the sublaminate.

n -RECOVERING OF LAMINA STRESSES

When the effective medium is used for stress analysis, the effective stresses {a} and strains {e} are obtained. The actual stresses and strains in a lamina are recovered as follows.

From the stress-strain relations {ak} = [Ck]{Ek} for the kth lamina in a sublaminate, we have

{ ~:} = {~::} = [::~ ::: ':::] {~} + [::: xz axz CSl CS2 CS6 Yxy CS3 C4S

k k k

C34 c3S] {Ezz} C45 Eyz

C55 k Yxz k

(7)

475

in which azz • ayz. au,. £Xx. Evv. and YXy are calculated from the analysis of the effective medium. The out-of-plane strains can be obtained from Eq. (7) as

{EzZ} [C33 Eyz = C43

'Yxz cS3 k

c34 c3S ]

-I

C44 C45

c4S css k {::} - ::: ::: ::] {:;}

axz C51 CS2 CS6 k 'Yxy k

(8)

Frcim'Eq: (1). the other three strain components E~x. E~. and 'fxy are seen equal_to the c~rresponding effective strains. Thus. {Ek} is completely determined from {O"} and {E}.

UI· GLOBAL·LOCAL ANALYSIS In the present global-local analysis, the thick laminate is divided into a global

domain and a local region. The global domain is modeled by effective properties, and in the local region the original material properties are retained. The actual lamina stresses in the global domain are recovered using the procedure discussed in the previous section.

A nine-noded isoparametric finite el«ment is used to model the global domain as well as the local region. The stiffness matrix for the finite element in the global domain is formulated with the effective properties. while the actual lamina properties are used in the local region. In addition, a full scale finite element model is used to analyze the original laminate. The result of the full scale finite element model is used for comparison purposes.

In formulating the finite element for the effective medium in the global r~gion, care must be exercised in modeling the free edge. Since the effective stresses O"u are averaged quantities ( over the sublaminate), stresses may exist in the indiVIdual laminas at the free edge. To ensure the exact traction free boundary condition. tractions of opposite sign must be added to cancel out these stresses. This results in local moments which are self-balanced over the thickness of symmetric laminates, but not for unsymmetric laminates. This added traction is given by

tj = -O"jknk (9)

where O"jk are the stresses in the lamina recovered from the effective stresses and strains. and nk is the outward unit normal vector to the free surface.

The nodal forces in a finite element (for the effective medium) corresponding to the traction given by Eq. (9) are calculated using

fjj = -f Njtjdf (10) r

where Ni is t~e shape function at node i, f is the boundary contour of the finite element, and Pi is the j-component of the nodal force at node i.

IV • NUMERICAL RESULTS The [Ol9OhOs and [05/90S]205 laminates are used for numerical examples. The

lamina elastic moduli are assumed as

476

El = 138 GPa, E2 = E3 = 10 GPa, G12 = G23 = G13 = 5.5 GPa,

V12 = v23 = v13 = 0.29

The effective moduli are calculated according to Eqns. (3-6). We have

Ex = Ey = 74.4 GPa, Hz = 10.7 GPa,Gxy = Gyz = Gxz = 5.5 GPa,

vxy = 0.039, VYZ = Vxz = 0.32

Two examples are illustrated.

4.1. Free edge stresses

(11)

(12)

A [O/90hOs laminate subjected to a constant strain Ex = Eo = 0.001 with dimension b = 4h = 1 cm is analyzed. The displacement field in this pseudo three­dimensional problem is given by

u = Eox+ U (y,z)

v = V (y,z)

w=W(y,z)

where u, v, ware displacements in x, y, z directions, respectively.

(13)

Figure 1 shows the laminate dimensions, coordinate system, and the global and local regions. Due to symmetry, only a quadrant of the cross-section is taken for consideration. Three finite element models are used. The all local solution is obtained using full scale finite elements tomodel the actual laminas. The other two solutions correspond to the use of one and two sublaminates, respectively, as the local region. .

Figure 2 presents the interlaminar normal stresses Ozz at midplane of the laminate. It is evident that these three methods yield identical stresses in the far field (away from free edge). Near the free edge, the use of two sublaminates as local region seems to result in a solution which converges to the full scale finite element solution.

4.2. Delamination cracks Delamination is one of the major modes of failure in composite laminates. The

stress field near the delamination crack and strain energy release rate are of interest to researchers. In this example, a [05/905hOs laminate containing an interlaminar crack located between the 900 and 00 laminas in the sublaminate just above the midplane is considered, see Fig. 3. The laminate is assumed to have infinite length in the x­direction and is subjected to a uniform stress 0 0 = 414 KPa on both top and bottom surfaces as shown in Fig. 3.

Two finite element models are employed; i.e., the global-local model, and the global model. In the global-local model, the local region contains 4 to 16 sublaminates depending on the crack size. In the global model, the whole laminate is represented by the equivalent continuum with the effective moduli given by Eq. (12).

Figure 3 presents the in-plane stress Oyy distribution along the thickness direction at y/a = 1.0156. The local stress recovered from the effective modulus theory is found to agree with the global-local solution quite well, even at locations near the tip.

The strain energy release rate is calculated using the crack closure method [7]. The strain energy release rates for Mode I and Mode II are given by

1 a+t.a

Gr = - lim f O"zz(y;a)~w(y;a+~a)dy 2~a t.a40 a

1 a+t.a

GIl = - lim f O"yz(y;a)~v(y;a+~a)dy 2~a t.a40 a

477

(14)

respectively. In Eq. (14), ~w and ~v are the near tip relative displacement of the two crack surfaces after fictitious crack extension of ~a. These integrals can be approximated using nodal forces and displacements from the finite element solution. It should be noted that for an interfacial crack between two dissimilar materials, G1 and Grr do not exist individually [8]. However, the total strain energy release rate G = Gr + Grr is well defined.

Table 1 lists the total strain energy release rates for different crack sizes. Comparison of the solutions of the two models indicates that when the crack size is approximately equal to or larger than the thickness of the sublaminate, the effective modulus alone is adequate for evaluating the total strain energy release rate of delamination cracks in composite laminates.

v -CONCLUSIONS

It has been demonstrated that the effective modulus model can be used with significant accuracy to solve free edge stress and delamination crack problems in thick laminates. For delamination cracks located in the interior of the laminate, the effective modulus model can be used alone if the crack size is equal to or larger than the thickness of the sublaminate. For free edge stresses, a global-local procedure is needed to obtain accurate stresses near the free edge.

Acknowledgement This work was partially supported by the Office of Naval Research under

Contract No. NOOOI4-84-k-0554 with Purdue University. Dr. Y. Rajapakse was the technical monitor.

VI-REFERENCES

1. G.W. Postma, "Wave Propagation in Stratified Medium," Geophysics, Vol.20 (1955) 780-806

2. N.J. Pagano, "Exact Moduli of Anisotropic Laminates," in Composite Materials (Editors: L.J. Broutrnan and R.H. Krock), Vol. 2, Mechanics of Composite Materials, (Editor: G.P. Sendeckyi), (1974) 23-45 Academic Press, New York, NY

3. C.T. Sun and S. Li, "Three-Dimensional Effective Elastic Constants for Thick Laminates," J. Composite Materials, Vol. 22, No.7 (1988) 629-639

4. R.B. Enie and R.R. Rizzo, "Three-Dimensional Laminate Moduli," J. of Composite Materials, Vol. 14 (1970) 150-154

5. N.J. Pagano and S.R. Soni, "Global-local Laminate Variational Model," lnt. J. Solids Structures, Vol. 19, No.3 (1983) 207-228

6. A.S.D. Wang, "Calculation of Edge Stresses in Multi-layer Laminates by Sub­Structuring," J. Composite Materials, Vol. 12 (1978) 76-83

478

7. E.F. Rybicki and M.F. Kanninen, "A Finite Element Calculation of Stress Intensity Factors by a Modified Crack Closure Integral," Engng. Fracture Mech., Vol. 9 (1977) 931-938

8. C.T. Sun and C.J. Jih, "On Strain Energy Release Rates for Interfacial Cracks in Bi-Material Media," Engng. Fracture Mech., Vol. 28, No.1 (1987) 13-20

Table 1. Interlaminar Crack Energy Release rates GE2/(cr~a)

aid 0.1 0.5 2.0 8.0 G (global-local) 6.253E-2 5.913E-2 6. 110E-2 7.396E-2

G (global) 5.950E-2 5.977E-2 6.005E-2 7.333E-2

difference 4.85% 1.04% 1.71% 0.85%

d = thickness of sublaminate

z y

global h

b

Fig. 1 - Global and local regions in a quadrant of the laminate cross-section

O,OIl...--------------------,

[0/901Ios b == 4h == 1.02 cm

0_06 __ all local

----.:i 0 ,04 ____ 2 sublaminates local + global

- -- 1 ublaminate local + global N

~

J 0.02

~ - ~ 0_00 b=====----...:s:~~=::~

-0.02 1--~--~-~~----~--'I, (Iil 0,70 0.75 O,Il0 n,II5 O,UO U.%

ylb

Fig. 2 - Interlaminar nonnal stress aU. along midplane

8 ' Z

6 0 0

Y

2 0 0

~ 2b

~ 0

lOs/90s 120< -2 b == h = 2.54 cm aid =8

·4 - global-local

- -- all global (recovered stress)

-6 d = thickness of sublarninate t = lamina thickness == 0.635 mm

- ~ .. 1.0 00 -1.0 ,0 12.0 16,0 20,0

Oyy/oo

479

Fig. 3 - InpIane stress Oyy distribution along thickness direction at y/a == 1.0156

ABSTRACT

MICROFRACTOGRAPHY OF CARBON FIBRE· REINFORCED BISMALEIMIDES

G. MAIER, M. WOHLEKE, P. VETESNIK, J. KUNZ

Wehrwissenschaftliches /nstitut fOr Materia/untersuchungen Landshuter Strasse 70 - 8058 ERD/NG - West Germany

For a well-founded failure analysis of parts made of composites we are still lacking sufficient information, especially in the case of mate­rials with a brittle matrix. This study therefore was aimed at investi­gating the influence of loading conditions on the fracture surface struc­ture of a carbon fibre-reinforced bismaleimide. The results have revealed that also in the case of brittle matrix materials an allocation of fracture surface structure characteristics to conditions existing during damage development is possible. Fracture surface structures in carbon fibre-reinforced blsmaleimides partly differ from those in carbon fibre-reinforced epoxy resins.

INTRODUCTION

In -service failures of components can never be precluded comple­tely. In many cases, especially if injuries or high material damage have occured, defect components wHl probably be examined for characteristics that allow determination of the cause of failure. Experience gained with conventional materials has shown that such characteristics can often be identified on fracture surfaces by using scanning electron microscope (SEM) techniques. To determine and evaluate these characteristics, a fundamental knowledge of the influence of loading and environmental conditions on the fracture surface structure is required. For composites, however, such knowledge is only available to a limited extent. This is especially true for composites using matrix materials suitable for high temperatures (e.g. polyimides, blsmaleimides). As a rule these matrix materials exhibit a lower elongation at break and thus a more brittle behaviour than the currently used epoxy resins for low and standard service temperatures.

482

I - EXPERIMENTAL

In a multi-directional laminate (SIGRI CI 1020) of stacking sequence (Oz/±45/0z/±45/90)s fracture surfaces were produced using defined test conditions. For this purpose specimens (width = 15 mm, length = 50-100 mm) were quasi-statically and cyclically (constant amplitude, R = 0.1. frequency = 10 Hz) loaded in a four-point flexure test. Preliminary tests have revealed that best results in controlled crack propagation (delami­nation) between two plies of the laminate used can be achieved with notched specimens having a notch depth of 4 plies. The loading-mode (mixed-mode or pure mode II) was varied by changing the notch position in the flexure test fixture. A mixed-mode loading (modes I & II) could be achieved by positioning the notch within the tensile stress area of the specimen. By positioning the notch within the compression stress area, mode I loading could be almost completely suppressed and thus a pure mode II loading achieved. These mechanical loading tests were performed at temperatures of 23 'C and 170 ·C. The fracture surface structures produced were evaluated by means of a SEM.

II - RESULTS

2.1. General

As already found on carbon fibre-reinforced epoxy resins (CFREP) /1-4/, the material investigated also revealed the structure types of hackles, river markings and textured microflow, but in a modified form. On both cyclically and quasi -statically loaded specimens the direction of the hackle formation could be attributed to that of the damage propaga­tion direction. Contrary to CFREP /4/, no fatigue striations could be ob­served on cyclically loaded specimen under the test conditions selected.

Particular attention was paid to the characterization of the border areas of the fracture surface. Experience has shown that many cracks of a component are initiated in these areas and thus knowledge of these structural characteristics would be of great value for later damage inve­stigations. Furthermore, even with other matrix materials, this area of the fracture surface has been hardly, if at· all covered in existing lite­rature.

2.2. Influence of loading mode at room temperature

In the case of mixed-mode loading river markings could be identi­fied as being the dominating fracture surface structure. Orientation of these river markings indicates a different crack propagation direction in border areas and at the specimen's centre (Fig. 1 and 2). Separation of plies (delamination) occured in the centre of the specimen on fibres of o '-plies; in border areas on fibres of both 45 '- and 0 '-plies. This resulted in higher and more narrow resin bridges in the centre than in the border areas of the specimen. Fracture surface structures in opposite border areas also showed a marked difference. In the centre of the spe­cimen, additional hackles occured, which, when compared with those

483

produced in CFREP /3-4/, showed a smaller slope angle and resembled a rounded sawtooth more than a comb tooth. These hackles point in the macroscopically induced crack propagation direction. As far as visible, these hackles originate in the crossing points of the 0 .- and 46 ·-ply fibres. At these crossing points textured micro flow structures form and then terminate in these hackles.

Under pure mode II loading the fracture surface structure appears to be less uniform than with mixed-mode loading (Fig. 4). Hackles, in particu­lar, are predominantly sharp-edged and, contrary to those observed on CFREP /3-4/, are irregular in shape. However, they then resemble more the shape of a comb tooth than that of a sawtooth (mixed mode) and are pointing in a direction opposite to that of the macroscopically induced crack propagation. This can be explained by the shear stress between the layers which, as compared with the mixed-mode loading, acts in the opposite direction. Border region and centre fracture surface structure of the specimen dif­fer less than in mixed-mode loading.

2.3. Influence of temperature increase to 170 ·C

An increase in temperature from 23 ·C to 170 ·C results in a mar­ked change of the fracture surface structure (Fig. 6). Hackles now show a more rounded and pronounced form and, especially with the mixed-mode loading, are much wider than on specimens tested at 23 ·C. Hackles also occur more . .frequently in the border areas of the specimen. In the centre of the specimen fibres of the 46 ·-ply are more frequently visible. Fur­thermore, in the fracture surfaces produced with mode II loading re­latively high and smooth resin bridges can also be frequently observed (Fig. 6). These two latter fracture surface changes indicate an increase in adhesive fallure of the fibre/matrix bond.

REFERENCES

1. G.E. Morris ASTM STP 696,(1979), 223-273

2. R.E. Robertson, V.E. Midroin J.o. Materials Science 20(1986) 2801-2806

3. B. Smith, R. Grove, T. Munns in "Failure Analysis of CompOSite Structure Materials" Material Lab., Wright Patterson Air Force Base, Dayton(OH,USA) AFWAL-TR-86-4033 , May 1986

4. H.E. Franz Zeitschrift filr Werkstofftechnik 16 (1986), 321-391

484

o river markings

o bed of o '-ply fibres

6 hackles

Figure 1: River markings in the centre of the specimen; static mixed mode loading at 23 'C; (X2100)

direction of crack propagation

0 river markings

6 fibres of 45 '-ply

0 bed of o '-ply fibres

/ Figure 2: River markings at the border area of the specimen;

direction of crack propagation approximately corresponds to the direction of the 45 '-ply fibres; static mixed mode loading at 23 'C; (X2100)

direction of crack propagation

0

6

0

485

river markings

fibres of 45 '-ply

bed of o '-ply fibres

direction of crack propagation

Figure 3: River markings at the border area of the specimen; direction of crack propagation is approximately perpendicular to the direction of the 45 '-ply fibres; static mixed mode loading at 23 'C; (X2100)

0 comb tooth shaped hackles

6 fibres of 45 '-ply

0 bed of o '-ply fibres

direction ---I"~ of crack

propagation

Figure 4: Sharped edged and irregular structure in the centre of the specimen; static mode II loading at 23 'C; (X2100)

486

0 round und wide hackles

0:, fibres of 45 '-ply

0 bed of o '-ply fibres

direction ----'..-- of crack

propagation

Figure 5: Rounded and wide hackles in the centre of the specimen; static mixed mode loading at 170 'C; (X2100)

o smooth resin bridges

/\ fibres of W 45 '-ply

o bed of o '-ply fibres

direction -----'" of crack

propagation

Figure 6: Smooth resin bridges; exposed 45 '-ply in the centre of the specimen; static mode II loading at 170 'C; (X2100)

POROSITY IN ADVANCED COMPOSITE MATERIALS: ITS EVALUATION AND EFFECTS ON PERFORMANCES

A. GUEMES, MA MOLINA-COBOS, R. GONZALEZ-DiAl'

Abstract

ETSt Aeronauticos 'EUIT Aeronautica

Ciudad Universitaria - 28040 MADRID - Spain

Particle reinfored MMC's are interesting in application in arduous enviroments. Advantages of these composites are increased modulus, strength, high temperature properties and wear restistance. The thermal expansion is reduced. In this paper the microstructures and properties of composites with different particle additions, eg SiC, TiB2' Ti(C,N), AlN and Al203-platelets produced by powder metallurgy techniques are dicussed.

I. Introduction

Magnesium base materials are gaInIng in importance due to their low density for aerospace, space, military and automobile appli­cations. The ever increasing demands on these materials in recent years led to the development of high performance materials. Conventional material developments using precipitation and solid solution hardening and grain refinement was unable to eliminate some of the disadvantages of magnesium alloys as for example the low modulus, poor wear resistance, poor high temperature strength and high coefficient of thermal expansion. Particle and fibre reinforced magnesium matrix composites can, by suitable selection of matrix and additives, exhibit a combination of metallic and ceramic properties. Such a property profile opens the door to new applications such as bearing materials, pistons, gudgeon pins etc. The following methods can be used to manufacture particle rein­forced magnesium materials: casting by stirring in the particles (P. K. Rohatagi et al./l/) or by powder metallurgical techniques

488

(R. L. Trumper /2/, L. Erich /3/). Problems in the first method arise due to the reaction between the fibres and the matrix melt and also in the dispersion of the particles. In the powder metal­lurgical techniques eg consolidation by extrusion, it is possible to largely avoid chemical reactions. There are, however, greater problems in obtaining a satisfactory dispersion of the particles. A homogeneous dispersion is only possible if specia.l dispersion methods are employed. The preparation and properties of particle reinforced magnesium materials are discussed below. The develop­ment of microstructure and the properties of a AZ91 with SiC, TiB2' AIN, Ti (C, N)50:50 and A120~-platelets are presented and compared. The matrix was prepareu in form of powder by machining and attrition.

II. Production of Composites

The AZ91 powder was prepared by Mimeta S. A. Lausanne using mecha­nical methods. The fine powder with a maximum particle size of 63 vm was seperated out and mixed with the additive. The additive was either SiC-, AIN-particles or A1 203-platelets with a mean size of 6 vm. Table 1 gives the chemical composition of the powder.

~A~l~ _____ Z~n~r-~~Mn~.~ __ ~S~i~ __ ~C~u~ Be 9,5 % 1,0 % 0,30 % 0,3 % 0,2 % --15 ppm max.

Fig. 2 shows the size distribution of the powder used. The physical properties of the additives and the d~O values are collated in Table 2. Dispersion was carried out dry in a mixer followed by mechanical agglomeration in 8. ball mill.

particle SiC TiB2 AIN Ti(C,N) A1203-Platelets crystal type hex hex hex hex diameter (Vm) 6.5 5-7 5.0 6 6.2 density (g/cm 3 ) 3.21 4.51 3.26 5.18 3.90 Mohs hardness (max.13) 9.7 9.5 7.0 9.0 coeffcient of_5h~lmal 4.7 4.6- 5.5 8.4 4.6 expansion (10 K ) 6.4

Table 2: Properties of particles added.

The flow diagram, Fig. 5, shows the production route of particle reinforced composites. The compact materials were produced by powder metallurgical techniques. Extrusion was used to consolidate the powder.

III. Microstructures and Properties of Composites

The microstructures of the materials after consolidation are sho~~ in Fig. 4. With platelets, alignment in the direction of extrusion is possible (J. A. Black /4/). On the other hand, particle rein­forced material shqwed no difference between and the longitudinal and transverse directions. The mechanical properties of the particle reinforced composites with magnesium matrix are shown in

489

Fig. 6. With an addition of 15 vol.-% of different particle reinforcements jncreasing strength and hardness was observed. Only the composites with A1N-particles showed a decreased strength. Wetting problems, chemical reactions and the powder size distri­bution are reasons for this behaviour. The youngs's modulus, Fig. 7, and wear resistance, Fig. 9, are increased with the particle addition. An decreasing thermal expansions coefficient, Fig. 8, was observed.

IV. Conclusion

The results show that is possible to produce particle reinforced magnesium using P/M-techniques with interesting properties. Particle reinforced composites represent one of the most inexpen­sive promising materials for automotive applications and have been sucessfully tried out as bearings, pistons and cylinder liners. Composites offer an improvement of the mechanical properties and wear restistance, but there are many problems during production and further working on these materials eg the reactivity of magnesium, the disperSion of particles depending on the process conditions, and wetting problems. It can be seen that not every addition material is suitable for particle reinforcement. Never­theless homogenous distribution and increasing mechanical proper­ties, wear resistance and reduction of the thermal expansion coefficient are possible, using appropriate P/M techniques.

V. References

/1/ P. K. Rohatagi, R. Asthana, S. Dee; Solidifiation, structures and properties of cast metal-ceramic particle composites, Int. Metals Review (1986) 31, 3, 115 - 136.

/2/ R. L. Trumper; Metal Matrix Composites - Applications and Prospects, Metals and Materials (1987), 662 - 667.

/3/ L. Erich; Metal-Matrix Composites, The Int. J. of Powder Metallurgy (1987) 23, 1, 45 - 54.

/4/ J. A. Black; Shaping Reinforcements for Composites, Advanced Mat. & Processes, Met. Process. 3 (1988) 51.

Fig. 1: SEM of magnesium powder AZ91.

logorithrnits Gaussian distribJtion

30

d ..... tor(ml

Fig. 2: Size distribution of matrix alloy AZ 91.

490

Fig. 3: Particle morphologie of additives: a. SiC d. AIN b. TiB c. Ti(t,N)50:50

Fig. 4: Microstructure of composites after extrusion. a. AZ91 + 15 vol.-% SiC b. AZ91 + 15 vol.-% TiE c. AZ91 + 15 vol.-% Ti(N,C)50:50 d. AZ91 + 15 vol.-% AIN e. AZ91 + 15 vol.-% A1203-Platelets

light Ntal alloys AI-,!1g- bas.

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492

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o A.191lulrus.,.,1 81HV10 o A.191·S.c IlSHVIO • A291'T,62 14OHVIO • A29I'T'(( ,N) 150HVlO • A291· ... IN 1l0HV 10 • AZ91 • .o.1 203 - plalelets

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MATRIX CRACKING IN CROSS PLIED THERMOSETTING AND THERMOPLASTIC COMPOSITES DURING

MONOTONIC TENSILE LOADING

R. VAN DAELE, I. VERPOEST, P. de MEESTER

Katholieke Universiteit Leuven De Croylaan 2 - 3030 HEVERLEE - Belgium

ABSTRACT

The evolution of matrix cracks during incremental monotonic loading. was followed using in situ radiography and was documented as a function of strain. Three commercial available matrices with a matrix failure strain ranging from low to very high were selected. From these tests can be concluded that the matrix properties greatly influence initiation and propagation of matrix cracks.

INTRODUCTION

One of the limitations of advanced composites. consisting of aligned continuous fibres in a polymer matrix. is that the strength and stiffness are quite low at substantial deviation from the fibre direction. As a consequence. multiply laminates with different fibre orientations are used. This has as result that the layers oriented normal to the principal loading direction are apt to fail at low stresses. Damage tolerance of composite materials can be determined for different loading conditions. The knowledge of the behaviour of the material under cyclic (fatigue) and dynamic (impact) loading are of utmost importance. However. to understand the basic mechanisms of damage development. tensile tests on crossply (°2 , 9°2) s laminates can be very useful. provided the damage can be monitored accurately. This paper will treat the detection of delaminations and matrix cracks using penetrant enhanced in situ X-ray radiography/l/ and the

494

influence of material and processing conditions on transverse cracking.

TEST PROCEDURE

A monotonic tensile test was carried out with small strain increments (0.05%). To document matrix crack initiation and growth, penetrant enhanced in situ radiography was used. The penetration is realized by submersion, while the specimen is strained, in order to have full and reproducible penetration. Figure 1 shows the experimental set-up with X-ray tube, specimen and penetration device. Comparative studies have shown that submersion penetration gives superior results, compared to conventional edge penetration.

A series of radiographs were taken during incremental monotonic tensile loading and the number of cracks and the total crack length were counted. There were some difficulties with APC-2 in determining correctly the number of cracks, due to the presence of small long voids. In the X-ray images, these voids produced features which were rather similar to those of matrix cracks. To overcome the problem of discriminating voids from matrix cracks, the APC-2 specimens were pre impregnated with a penetrant of lower X-ray absorption. The void distribution was then recorded using X-ray radiography. The incremental tensile testing was performed using a penetrant with a high X-ray absorption (Di jodo methane, DIM) and X-ray conditions which didn't reveal the voids filled with the penetrant of lower X­ray absorption. The resulting X-rays were compared to the damage visible at the edges as a control measure.

MATERIALS

The initiation and propagation of matrix cracks is influenced by a number of parameters. First, we have the material properties itself like G1C and failure strain (Ef ), plastic deformability of the matrix and the properties of fibre-matrix interface. Secondly, the processing parameters change the laminate resistance to matrix cracking by altering the residual stresses, crystallinity (thermoplastics), presence of voids, and degree of cure(thermosets). Thirdly, fibre volume fraction, the local geometry and laminate lay­up will produce stress and strain concentrations which will largely influence matrix cracking.

Three materials were selected on basis of matrix properties like strain to failure and fracture toughness (table 1).

The carbon epoxy prepregs were laminated and cured in a press clave according to the manufactures prescribed cure cycle . The carbon-PEEK (APC-2) laminates were produced by ICI. All mentioned laminates are symmetrical (02 902)s crossplies.

RESULTS AND DISCUSSION

When looking at the crack growth of IM6/6376, a rather sudden increase in the number of matrix cracks is found. The strain at which

495

severe cracking starts is low compared with the matrix failure strain. The matrix cracks initiate and grow from 0.4% strain onwards to reach some sort of saturation level. In APC-2, matrix cracks initiate but do not grow deep into the material, even at relatively high strains. Several factors are responsible for the different transverse cracking.

First, there is the residual strain due to the differences in thermal expansion of the matrix and the fibres. The residual stresses parallel to the loading direction in the 90· layers are calculated for the three systems, using for the epoxies an estimation of the stress free temperature based on the curing temperature and resin glass transition temperature Tg. The residual stress and strain is given in table 2. The magnitude of the residual stresses and strains are high compared to the transverse strength and transverse failure strain of unidirectional laminates.

The second contributing factor is the stress and strain concentration in the matrix near the fibres due to the different elastic properties of the fibres and the matrix. The larger the difference in elastic properties, the larger the stresses and strains at the interface. Skudra/2/ calculated the stresses at the fibre-matrix inte~face using a diamond array of fibres. This gave for carbon epoxy a stress concentration factor of about 1.4 for carbon epoxy and 1.8 for glass epoxy. These stress concentration factors depend on the stiffness ratio of the fibres and matrix. Since the transverse modulus of the different carbon fibres used, do not differ much, and the matrices have about the same stiffness, the stress concentration at the interface will be nearly the same for the three composites involved in this research.

Skudra also showed that for carbon epoxy, local variations in fibre density will not affect the strain concentration factor substantial­ly.

The expected advantages of the high failure strain of APC-2 are not realised in the transverse properties of a unidirectional laminate. The transverse stress and strain have indeed improved but not to their possibilities when one regards the high matrix failure strain and the relatively low strain magnification involved.

Even though matrix failure strain, thermal stresses and strain concentrations cannot explain the unexpected fast initiation of matrix cracks in the ductile polymers, they can explain the propagation of matrix cracks. The two matrices with plastic deformability (PEEK, 920-epoxy), high fracture toughness and high failure strain show clearly a better performance when looking to the crack growth.

Further research however, is needed to get better insight in initiation and propagation of transverse cracking in carbon reinforced plastics.

496

CONCLUSION

The controlling resin properties contributing to the transverse failure process still remain unclear due to the extremely complicated stress state in the heterogeneous fibre-matrix structure.

However, it is apparent that the matrix properties play an important role in the strain level at which transverse cracking starts and the evolution of these cracks. Brittle matrices with small strain to failure know fast crack extension once the crack is initiated, while more ductile matrices are characterised by slow crack extension.

ACKNOWLEDGEMENT

The preceding text represents research results of the Belgian programme on interuniversity attraction poles initiated by the Belgian state - Prime Minister's office - Science Policy programming. The scientific responsibility is assumed by it's authors.

REFERENCES

1 R. Van Daele, I. Verpoest, P. De Meester, I situ radiography as a means for calibrating acoustic emission, Composites Evaluation, Proc of the 2nd Int. conf on Testing, Evaluation and Quality control of Composites, TECQ-87, Univ of Surrey, Guildford, 22-24 sept 1987, p 117-128

2 SKUDRA, A.M. Micromechanisms of failure of reinforced plastics, Handbook of composites, vol 3, 1985,pl-69.

3 JERONIMIDIS, G., PARKYN, T.A., Residual stresses in carbon fibre thermoplastic matrix laminates, J. of Comp Materials, vol 22, 1988, P 401.

strain to fracture failure Energy % J/m2

Fibredux 6376# 3.1 432 Carbon epoxy Fibredux 9201 >7 Carbon epoxy Polyether- 30-150* 4800 etherketone+ Carbon Thermoplastic

* depending on the degree of crystallinity I Ciba Geigy + I.C.I. ,APC-2

Table 1 : matrix properties

modulus

GPa

3.6

3.7

material composite UD 90° fracture stress free failure energy** temp*** stress J/m2 K MFa

Fibredux 6376 0. 47 180 56 Fibredux 920 0.9 100 50 APC-2 1.4 310 80

** Double cantilever beam test, initiation values ***loss tangent maximum

Table 2 Laminate properties

residual transverse expected thermal thermal failure strain

material stress strain Er strain Ef Ef-Er MFa % % %

Fibredux 6376 47 0.4 0.50 0.1

Fibredux 920 25 0.3 0. 62 0.3

APC-2 40 /3/ 0.45 0.89 0.34

Table 3 Mechanical properties of the selected laminates.

figure 1 : in situ radiography as it is mounted on the tensile testing machine

497

measured strain 1st cracks %

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498

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Figure 2a : Total crack length evolution Figure 2b : Total crack length evolution during

during monotonic tensile loading in APC-2 monotonic tensile loading in Fibredux 6376

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.t 12 ., en c •

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Figure 2c : total crack length evolution

during monotonic tensile loading in Fibredux 920

ENDOMMAGEMENT ET FATIGUE DAMAGE AND FATIGUE

Chairman: Mr N. SPRECHER Owens-Corning Fiberglas Europe

DAMAGE DEVELOPMENT IN CARBON FIBRE REINFORCED COMPOSITE LAMINATES

UNDER COMPRESSIVE STATIC AND FATIGUE LOADING

K. SCHULTE, J.J. MASSON

DFVLR - Linder HOhe, Postbox 906058 - 5000 KOLN 90 - West Germany

ABSTRACT

Composite materials reinforced with continuous fibres offer superior tensile properties under static as well as fatigue loading. However, under compressive loading the properties seem to be not as pronounced. Buckling of the test coupons falsifies the results in a way that not the real material properties, but specimen geometry, laminate lay-up, fibre diameter or even the unsupported length of the test coupons determine the results achieved /1-3/. While under tensile loading the composite properties are mainly dependent on the fibres, under compression the matrix gaines more and more influence.

INTRODUCTION

The present investigation focuses on a comparative study of the compression testing of various laminates ([~s), [90 s ), [9016 ], [02,902,02,902~ ) .. As.buckling is a maJor problem under compressIve IOaa1~g 1t 1S prevented in the mechanical tests by using an antibuckling guide. If the unsupported length L of the test coupons was varied, the Euler equations desc¥ibing buckling overestimate the buckling strength of a composite (compare Fig. 1). Only if the unsupported length of the test coupons, tested in this investigation, was about 12,5 mm or below acceptable com­pressive data were achieved.

In this investigation special attention has been given to experimentally determine the compression stresses of the

502

neat matrix resin. A special test fixture (Fig. 2) was de­signed. The dimension of the neat resin specimen was 10 x 15 x 4 mm. All static tests were performed at room tempera­ture, but at different cross-head speads. Fatigue behavior was studied by cyclic loading at relatively high stress amplitudes with a varying between 50 - 90 % of the static strength. Fatigue tests were performed at different stress ratios R (R = a . /0 ) of 0,1; -1; -0,55 and 10.

m~n max

Two different types of fibres were used in the study. The intermediate modulus fibre TSOO and the high tenacity fibre T300, both in the modified bismaleimide resin system 5245c from BASF-Narmco. The fibres were produced by Toray, Japan. The basic mechanical properties of the fibres are given in Table 1.

Damage development was investigated by extensive de­structive and non-destructive test techniques, as light and scanning electron microscopy, specimen deply and X-ray radiography.

1. STATIC TESTS

In Fig. 3 are shown the tension and compression pro­perties of the neat resin and of the unidirectional lami­nate [90 S] and [90~6). The tensile stress for the neat resin is about S3 MPa, the elastic modulus 3,3 GPa and the fracture strain 2,9 %, according to the manufacturer. In the compressive tensile test, the yield stress of the neat matrix varies between 153 and 173 MPa dependent on the cross head spead, the yield strain between S % and 10 %, while the elastic modulus stays in the range between 3,2 and 3,9 GPa. Both, the neat resin and the transverse lami­nate achieve under compression higher values as in tension.

Fig. 4 shows the distribution of stresses around a fibre in a composite, which is loaded in transverse tension (a), respectively transverse compression (b) /4/. When loaded in tension the stress concentration at the fibre/ matrix boundary initiates early failure, with the result, that the tensile fracture stress of the transverse speci­men is lower as the fracture stress of the neat resin. Also the fracture strain is reduced, compared to the neat resin. The fibre reinforcement only increases the elastic modulus. Macroscopically compressive failure in an isotro­pic alloy occurs because of shear stresses on the planes 45° to the loading axis. This is in general also true for unidirectional composites, however, shear failure occurs in the planes which exhibit the fibre direction. Microme­chanically compressive failure in a [9016) laminate occurs, because tensile stresses are induced at the fibre matrix boundary transverse to the loading direction (Fig. 4 (b}). Failure initiates in tension at the fibre matrix boundary

503

leading to a highly irregular crack front along the fibre/ matrix boundaries (Fig. 5.2). But also matrix shear failure can be observed at certain locations (Fig. 5.1), where resin rich areas are present. failure is additionally ob­served as a combination of shear and compressive stresses (Fig. 5.3). The superior compressive stresses, when com­pared to the tensile stresses are related to the fact that:

- Under a compressive load the crack tends to shear in the 45° direction. The fibres support each other, respectively the irregular crack front hooks up (Fig. 5.2).

- Failure initiates under compressive load in transverse tension at the fibre matrix boundary (Fig. 4(b)).

Even the neat resin has higher strength in the compressive direction than in tensile direction.

One basic result of the tests of the unidirectional [0 ] and cross-ply laminates under static loading is that th~ laminates containing the T800 and T300 fibres do offer similar properties in compression, regardless the fact that the tensile properties of the laminates with the T800 fibres are superior to those with the T300 fibres /1/. The advantage of using the T800 fibre in compression loading is mainly related to their higher elastic modulus, which in tension and in compression exceeds that of the laminates with the T300 fibre.

2. FATIGUE TESTS

In Fig. 6 results of fatigue tests, all containing compressive load cycles, are summarized, where 10 . I, the maximum compressive stress in each load cycle is ~!§tted versus the number of cycles, N; and this for two different laminates. It is evident that pure compressive cycling results in an extremely flat 10 . I vs. N curve. Introduc­ing an additional portion of teW§~le loading significantly reduces the fatigue life (R = -1). This is more pronounced the higher the tensile portion in each load cycle is (R = -0.55); e.g., a specimen with the laminate ssacking sequence of [0 2 ,90 2 ,0 2 ,90 2 ] would not fail before 10 load cycles under tension-tens~on fatigue, when loaded with a maximum stress of about 750 MPa. However, it will fail far earlier, even when fatigued at a lower stress level by introducing additional compressive stresses. This clearly shows that there is a synergistic effect if tensile and compressive stresses are combined.

3. DAMAGE DEVELOPMENT

The compressive loading of a carbon/epoxy composite produces a somewhat different scenario from tensile load­ing. The major damage mode under static compression is

504

the kinking of fibres in the 0°-plies followed by delamina­tion and subsequent buckling of sublaminates. The results of experimental investigations show shear crippling as a possible failure mode. The shear crippling is a result of a kinking failure of the principle load carrying fibres in a laminate. It is characterized by a band of buckled, frac­tured fibres that has undergone both shearing and compres­sional deformation. The individual fibres in a kink band do not fail perpendic­ular to the fibre direction, as it is the case when fibres fail under tension load, but under some angle. Due to the compressive load the fibres are shearing off, sometimes under a quite low angle. This is shown in an SEM-micrograph in Fig. 7. A 0°-ply containing a kink band will move relative to the neighbouring out of axis plies. This brings out interply shear stresses following delamination between the failed 0°-ply and the angle ply. After delamination the sublami­nates are more susceptible to buckling than the original laminate, because of thinner thickness the momentum of inertia has decreased. The global buckling of the sublami­nates now leads to final failure.

Under compressive fatigue loading failure mainly oc­curs due to the onset of delamination which in the follow­ing load cycles allows parts of the laminate to buckle and initiates now, very rapidly within only few cycles, the final failure of the specimen by the loss of the specimen stability, as described above. In the case of additional tensile portions in each load cycle the typical damage pattern for tensile loading, as transverse and longitudinal cracking, develops. This damage pattern now allows to more easy initiate a delamination and now again the rapid final failure occurs.

ACKNOWLEDGEMENT

The authors gratefully acknowledge the support by the Deutsche Forschungsgemeinschaft (DFG) and thank the BASF company for supplying the prepreg system. REFERENCES

/1/ Masson J.J., Baron C. and Schulte K., Proc.: JNC 6, Sixiemes Journees, Nationales sur les composites, 11-13 Oct. 1988, Paris, 661-671.

/2/ Henrat P., in "Looking ahead for Materials and Processes" (J. de Bossu, G. Briens, P. Lissac, ed.) Amsterdam 1987, 389-399.

/3/ Lifshitz J.M., J. Composites Technology and Research 11 (1988) 100-106.

/4/ Newaz G.M., SAMPE Quarterly (1984) 20-26.

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Fig. 4:

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Stress distribution around a trans­verse fibre (90°) under tensile (a) and compressive (b) loading, after Newaz /4/.

3

The various damage mechanisms under transverse compressive loading in an unidirectional laminate.

480 T800 I 5245 [ MPa

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SEM-micrograph of broken QO-fibres under static compression. Matrix pyrolized.

507

ABSTRACT

DAMAGE TOLERANCE OF CARBON FIBRE REINFORCED PLASTIC SANDWICH PANELS

K. LEVIN

The Aeronautical Research Institute PO Box 11021 - 161 11 BROMM A - Sweden

This paper presents an experimental evaluation of static strength and damage growth resistance under spectrum loading of impacted carbon fibre reinforced epoxy sandwich panels. The program includes determination of the influence of impact energy on damage visibility, damage size and distribution, and strength losses with respect to different skin thicknesses and core densities. In the case of barely visible impact damage the susceptibility of sandwich panels to impacts is reflected in large strength reduction and furthermore in damage growth during fatigue loading. The visibility of damage during the fatigue loading has been subsequently reduced. Damage growth and failure also occurred for non-visual impact damages.

I INTRODUCTION

The use of carbon fibre reinforced plastics in aircraft components has increased rapidly in recent years. Of great concern is the susceptibility of laminated composites to impact damage. The sensitivity of plain laminated skin to impact can result in a substantial decrease in strength especially in compression loading [1-2). The problem arises since the strength can be reduced to the level of typical design strain used in aircraft structures. Furthermore the damage is not always visible from the front side which makes detection more difficult. Therefore, the application of carbon fibre reinforced plastics in primarily loaded structures necessitates consideration of the effect of impact damage to obtain sufficient confidence in safety of the structure. Moreover service experience with thin-skinned honeycomb composite structures has shown, (3), that they can be sensitive to low energy impacts, which raises the cost for maintenance, repair and replacement.

510

The object of this study was to determine the susceptibility of sandwich panels in terms of static and fatigue strength to damage as a result of impact and also to characterize the nature of the impact damage. The influence of impact energy on damage visibility, damage size and distribution, and strength reduction was evaluated for different skin thicknesses and core densities.

2. EXPERIMENTS 2.1 Specimen Configuration

The sandwich panels were fabricated with commercially available CFRP material and aluminium honeycomb cores. The skins were made from T300/914C in the form of unidirectional prepreg tape. The curing temperature in the autoclave was chosen to 190°C and held for two hours. No post-curing was performed. After curing the quality was examined using ultrasonic C-scanning. The two identical skins were bonded to the core with the adhesive FM300K. Two different skin thicknesses, 8-plies and 28-plies, were included in the test program. The stacking sequences of the two skins were (±45/0/90)s, 25/50/25 lay-up, and (±45/90/0/90/0/±45/90/0/90/0/±45)s, 29/42129 lay-up, respectively. All three types of core had the same cell size, 3.2 mm, althought the core density differed. The three different core densities were 72, 130 and 192 kg/m 3. The thicker skin was combined with all of the three different cores whereas the thinner skin was only combined with the lightest core. The dimension of the sandwich panels was 1000mm long by 160mm wide. The height of the panels was in all cases 50mm.

2.2 Impact and Mechanical Testing

In this program all the impact locations were subjected to low velocity impacts using a falling weight with a 30mm hemispherical tip guided in a tube. These conditions were chosen to simulate tool drops etc. possibly occurring during assembly and maintenance. The test panels were clamped between two steel plates containing a square cut-out, 140 by 140mm where the skins were impacted in the centre. No loads were applied to the test panels in connection with impacting. To find the impact energy required to obtain barely visible impact damage (BVID), i.e. the energy level where it is likely to find the damage during visual inspection, trial impacts were performed. Barely visible impact damage was defined as a 2mm deep dent in the front surface. After impacting the maximum dent depth was determined. Ultrasonic C-scan evaluation was used to determine the size of the impact damage for each impact location. For a number of panels sectioning was employed to evaluate the damage distribution through the thickness of the damaged skin. In addition several impacts of lower energy levels were tested to obtain a wide range of damage sizes and indentations.

For the initial and residual static strength assessment the panels were tested in four point bending with the impact damage introduced in the compression loaded skin. For the fatigue loading a symmetrical fin spectrum was applied, tentatively, for two life-times. One life-time corresponded to 3000 flight hours. A maximum compression load occurred once in every 100 flight hours. The test frequency did not exceed 1.5Hz. The subsequent delamination development was monitored by C-scanning evaluation at a number of times during the fatigue testing. The mechanical testing was exclusively conducted at room temperature under ambient conditions, as were the impact tests.

511

3. RESULTS AND DISCUSSION 3.1 Impact Damage Size and Visibility

The effect of impact energy on damage size is shown in Fig. 1. The damage size is specified as the extension in the width direction of the ultrasonic C-scanning recordings providing a planar picture of the impact damage. These recordings did not generally deviate from a perceived circular shape. Generally higher impact energy resulted in larger impact damage but the increase in the impact damage sizes became less at higher impact energies. Furthermore small impact energies did also cause large damage in the skin. This behaviour was observed for all configurations. Results on the thick-skinned specimen with low core density showed more sensitivity to impact i.e the damage sizes differed by a factor of 1.5 from those for higher core densities. This might be related to the fact that the low density core is less able to absorb energy during the impact. On the other hand the difference in the energy level to obtain a barely visible impact damage, i.e. a dent depth of 2mm, was less pronounced. In Fig. 2 the influence of impact energy on dent depth is presented. It was indicated that the dent depth is proportional for the thin-skinned specimens to the impact energy.

The impact damage distributions in the thick-skinned panels obtained from the cross-sections show a general increase in delamination size through the whole thickness when the impact energy increased, leading to increasing core crushing. Comparison of two sectioned panels with different cores impacted at the same impact energy level indicated that the deformation of the lighter core was more extensive. The extent of fibre fractures differs as well. A slightly different delamination distribution was also observed in the skin between high and low density core panels. In the stiffer configuration, i.e high density core panel, the distribution of the delaminations was more uniform than in the case of low density core, where the largest delaminations occurred in low lower part of the impacted skin. However the delamination sizes in the different ply interfaces were generally substantial larger in the low density core panels. For the thinned­skin panels with impact energies of 8J and 11 J, no significant difference was seen in delamination distribution and size.

3.2 Strength and Fatigue Test Results

Table 1 presents a summary of the strength data and recorded damage sizes. Data for thick-skinned panels show an effect of core density on static failure strain of the skin. The static failure strain values for the low density core panels were approximately 0.30% whereas medium density core panels gave values which were at least 20% higher. A significantly larger damage size as a result of the lower core density in thick-skinned panels seems to have as a consequence lower strength. Moreover BVID has a more adverse effect on strength in thick­skinned panels than in thin-skinned panels with the same core. However more impact energy is required to obtain BVID in thick-skinned panels. It has been observed in other works [5-6] that the distribution of internal damage sizes in an impact damage affects compression strength. Near-surface delaminations seem to degrade the strength more than other defects. In the present study, where the damage distribution was characterized by sectioning the result has not been sufficient fully to explain the differences obtained in the static failure strength.

Notable for the damage growth behaviour was that of the low maximum strain in the spectrum corresponding to 80% of the damaged static strength, 0.30% and 0.24% for the thin-skinned and the thick-skinned panels

512

respectively. There was an extensive growth and in some cases failure occurred under fatigue loading, see Table 1. In the thin-skinned panel, where both barely visual and non-visual impact damage, 11J and BJ, were tested, panels have failed during spectrum loading and and therefore these impact energies can be a substantial threat to a component. Non-destructive inspection must consequently be applied at high stress areas in the component. When the panels were able to sustain the fatigue loading, it was observed that, despite the substantial growth which by using the ultrasonic C-scanning testing has been indicated to occur near the surface, the impact damage seems not to impair the residual strength further. This might be due to the fact that failure is not necessarily caused by large multiple delaminations. Impacts with the lowest energy, 1J, on the thin-skinned panel did not result in damage growth. In the thick-skinned panels, the impact growth was connected with strength reduction. This might be due to the large damage area in relation to the panel width.

An important observation was the diminishing of dent depth as a result of spectrum loading for all panels and impact energies, Fig 3. In the first few flight hours, the dent depth decreased by 20 to 30%. The quick change in dent depth was not associated with damage growth, but resulted in decreasing visibility of the impact damage during further aircraft operations. The result does not indicate that the phenomenon is more pronounced for one of the two panel types. However for the lowest impact energy, 1J, only slight changes occurred during the spectrum loading.

4. CONCWSlONS

The low static strength response and the poor growth resistance under typical spectrum loading conditions clearly demonstrate the susceptibility to impacts of laminated CFRP sandwich panels. Data show that especially the thick-skinned panels are severely degraded by BVID. It is alarming that under the maximum load in the spectrum, corresponding to 80% of damaged static strength, damage growth did occur for both BVID and non-visual impact damage together with a quick decrease of visibility in the first few flight hours after the impact event. A significant influence of core density on damage size was observed demonstrating the high sensitivity to impacts of panels with light cores.

ACKNOWLEDGMENTS

Particular appreciation is expressed to B. ThOrnqvist and A. Linder for· their skilful experimental assistance. The author is also indebted to the Saab-Scania Aircraft Division and to the Swedish Defence Material Administration for providing panels and financial support respectively.

REFERa\CES

1. Levin, K., Proc. of ASC First Technical Conference, (1986), p.313 2. Demuts, E., Whitehead, R. S. and Deo, R. B., J. Composite Structure,4(1),

(1985),p.45 3. Whitehead, R., Presented t014th Symposium of ICAS(1987) 4. Rhodes, D. R., NASA TM 78719, (1978) 5. Byers, B. A., NASA-CR-159293, (1980) 6. Guynn, E,G., and O'Brien, T.K., 26th Structures, Structural Dynamics

and Materials Conference, AIAA (1986), p.187

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STATIC AND FATIGUE FRACTURE OF COMPOSITES IN COMPLEX STATE OF STRESS

D.PERREUX,C.OYTANA,D.VARCHON

Laboratoire de Mecanique Appliquee URA 279 - Universite de Franche Comle - 25000 BESANCON - France

ABSTRACT

Some results of static and fatigue tests in complex stress states of tension and torsion are shown. The static capacities of three glass/epoxy materials are compared. Finally, the sensitivity of the average stress in the form of isonumber cycles to failure curves are shown.

INTRODUCTION

The behaviour under complex loading, widely found in reality, gives useful information regarding choice of solutions to design a structure. Because of their heterogenous constitution, it is impossible to find the fracture and fatigue behaviour in the case of real loads from simple tests, such as flexion or tensile tests. The importance of experimental study in complex loads is obvious.

The present study concerns three kinds of glass/epoxy materials. The static and fatigue capacities under complex loads of tension and torsion are shown.

- MATERIALS, SPECIMENS AND TESTING METHOD

1.1. Materials and specimens

Filament winding is used to obtain three thin-walled, tube-shaped materials from the same components (glass E 1200 Tex, epoxy matrix DGEBA :Vf-0.6). This allows the fabrication of several kinds of reinforced structures. Unidirectional (UD,+4S),laminate (L,[+45/-45J 3 ) and hybrid structures, like woven fabric,(W, ±45) are studied.

516

1.2. Testing method

All tests are made under proportional loading. In other words, the specimen is constrained by a torsion torque and tension force leading to stress as' ay (fig 1) in such a way that during a test

~ = Q is a constant. as

Two kinds of fatigue tests are carried out:

test with average stress (UD, W)

a

I a y (t) - ~ (sin wt + 1) (1) 2

ay(t)/as(t) - c>

test without average stress

{ ay(t) - aymax sin wt

(2) a y (t)/as (t) - Q

II - STATIC RESULTS

(W)

Various failure criteria have been proposed [1, 2, 3, 4, 5J for anisotrope material. Tsai-Wu's equation allows in many cases to obtain a good correlation using few coefficients.

2.1. Failure criterion of UD

In the 1.2 axis (orthotropic axis) fig 1, Tsai-Wu's function is represented as follows:

with

(4)

a6 -

(3) and (4)

(5) flO y a y

a -'L

2

a -'L

2

lead to

+ ~O as

Tsai-Wu's function in (x, y) axis (fig 2)

+ flo yy 2 + flO

0y ss Os 2 + 2~~ ay as - 1

It may be noted that Tsai-wu's equation is in accordance with the experimental points

2.2. Failure criteria of Land W

In this case, the orthotropic axis of materials is the same as the specimen axis which gives th~ following criteria:

L, W (fig 2)

It may be observed that W exhibits a tensile strength higher than L but shear stress is the same for both cases. The entanglement of strands in W reduces propogation of the transverse cracking matrix. This breaking mechanism is predominant in the tensile test. On the other hand, in the shear test, breaking fibres controls rupture; in this case there is no difference between these two materials.

The first ply failure criterium of L has been calculated from results obtained on UD. It may be noted [6, 7J that its use in the design of a structure seems very conservative.

III - FATIGUE RESULTS

3.1. Analysis method

For any values a we determine the Wolher curves associated for one kind of test. It is possible to represent them by straight lines (7? for the number of cycles to rupture (NR) comprised between 10 and 105

(7) j a (a, NR) ~ A(a) + B(a) Log(4NR) ymax

a (a, NR) 0smax(a, NR) ~ ~y~m~a~x~ ____ __

a

Let NRO fixed NR. In the first quadrant (Oy,os) some points (Oymax(a, NRv),osmax(Q, NRo)) drawing an isonumber of cycle to failure curve (ICFC).

3.2. Tests with average stress (UD, W)

of stress space are obtained,

This type of test can following equation is proposed describe ICFC (fig 3,4)

be compared with static tests. The in connection with static criteria to

(8) j F' L + F' L + F' y y s s yy

Ly ~ 0ymax . T~ (NR)

Ls~Osmax T~(NR)

1 . Bi Log 4NR

where T~: T~ are functions of NR (fig 5) which transfer the static criterla allowing to take into account different rates of growth of failure mechanisms. Bi is a constant representing an homothety of the static criterium, taking into account the sensitivity to fatigue (B~ ~ 0.15 BUD ~ 0.2).

517

518

3.3. Test without average stress

Now ICFC (fig 6) show a form criteria. An analogy between material may be noticed. We also tests, the specimen breaks during cycle.

which cannot be comf~red with static ICFC and the yield ) criterium of

observe that in traction compression the compressive part of the fatigue

(*) in fact, criteria of loss linearity in stress strain curves

IV - CONCLUSION

This test in complex stress shows that with the same components, the static characteristics can be improved by the correct choice of their unification.

Finally, we also observe that without general laws of behaviour we cannot determine the multiaxial behaviour from conventional tests.

REFERENCES

1 - Hill R., A theorie of yielding and plastic flow of anisotropic metals, Proceding of Royal Society, A 193 (1948), 281-297

2 - Gol'denblat 1.1. ,Kopnov V.A., Strenght of glass-reinforced plastics in complex stress state, Polymer Mechanic, vol 1 (1965),54-59

3 - Tsai S.W., Wu E.M., A anisotropic materials, Journal of 58-80

general theory of strenght for composite materials, vol 5 (1971),

4 - Hashin Z., Failure criteria for unidirectional fiber composites, Journal of applied mechanics, vol 47 (1980) ,329-334

5 - Boehler J.P., Delafin M., Failure criteria for unidirectional fiber-reinforced composites under confining pressure, Mechanical behavior of behavior of anisotropic solids (editions du CNRS),(1982),449-470

-6 - Perreux D., Oytana C., Varchon D. and Atcholi E.K., Etude experimental de la rupture statique en contraintes complexes, Comptes­rendus des sixlemes journees nationales sur les composites ,Paris (1988), 554-564

-7 - Oytana C., Perreux D. criteria on multiaxial stress Poceeding ICCK ,(1989)

and Varchon D., Static and fatigue test probelms posed by prediction,

519

Materials Fy F. Fyy F .. FyS

UD 9.97 E-3 -2.06 E-2 1. 12 E-4 2.00 E-4 -1. 04 E-4

S 3.32 E-3 o. 5. 53 E-5 9.31 E-6 o.

C 1. 24 E-3 o. 3.97 E-5 9.75 E-6 o.

Table Stat[c parametres of materials

Fig.l - state of stress ()., '. f1PA )

HO _

*: W _ _ -;.......,

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• : U 0 --'+-I..,L/

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520

A )lR - lSOOO • )lR - 2500 • ~R - 150

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-~,

30· MPa

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I I

: ..

A ~R - 25000

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T--~ . :,.

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0; MFa

I , I

-I : ~J_ jLL~

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Fig.6 - ICFC of W (tests without average stress)

DAMAGE DEVELOPMENT IN CFRP AND ITS DETECTION

R. AOKI, J. HEYDUCK

DFVLR Pfaffenwaldring 38-40 - 7000 STUTTGART 80 - West Germany

ABSTRAcr CFRP specimens, containing an idealized discontinuity, were

subjected to static tension tests and an irrpact danaged CFRP stringer stiffened panel was compression fatigue loaded to show same of the capabilities of the ultrasonic-insitu method. H-scans obtained with this method show a feasible way to correlate stiffness degradation and ultrasonic attenuation measurements.

INTROoucrICN The effect of defects and danages on strength, stiffness, stabi­

lity and service life of CFRP structural parts is one of the Irost inportant problems affecting the use of structural composite materi­als. Part of an ongoing research program in the Institute for Struc­tures and Design is to characterize the load-induced danage devel­opnent in CFRP specimens and structural parts under static and fatigue loading. CFRP flat specimens, forseen with an idealized discontinuity, and an irrpact danaged stringer stiffened panel were the test objects. The evaluation of the danage progression was done with the ultrasonic US- insitu method, developed in our Institute. The results show the feasibility to detect minor changes in the material and to correlate them with the stiffness degradation.

1 - EXPERIMENTAL PROCEDURE The flat specimens prepared from T300/Code69 prepreg had differ­

ent stacking sequences, changes of the quasi-isotropic (OI±45/90)s lay-up, as shown in Tab. 1. Earlier investigations on CFRP laminates subjected to low velocity, low energy irrpact III exhibited the

522

presence of fiber breakage and delamination at the opposite surface of the irrpacted region. To simulate this, the dam3.ge was idealized by perforating the 3 or 4 outer layers of the specimen before curing. The hole diameter of the perforation was 4mn. This diameter was chosen as representative fiber dam3.ged region in specimens (lIlm thick) with invisible dam3.ges in the front side. After the curing process the periphery of the idealized discontinuity could be distin­guished as a shadow. All specimens were straight-sided coupons 36mm or wider and 90mm in gage section. Specimens were endtabbed with crossply glass/epoxy tabs. The thickness of the laminate was in all cases lIlm. The structural part, a stringer stiffened panel, with the lay-up (0/±45/90/±45/02) , was fabricated with T300/Code 69 fabric after the one-shot method dg;,eloped in the institute /2/. This method allows to cure the stringers and the panel in one autoclave cycle, avoiding the problems of joining between stringers and the panel. The test panel in this investigation was 90x12Omm in size with an overall thickness of 2mm. The irrpact dam3.ges were introduced, in a drop weight facility /1/, in different locations of the panel, e.g. between the stringers, beside a stringer, and just over the other stringer, Fig. 1. The irrpact parameters were m=lkg, E=2,75J, hemispherical irrpactor with lOmm diameter. No dam3.ge at the front surface was visible. Static and fatigue tests were performed on a Schenk servo-controlled hydraulic test maschine. Fatigue tests were at constant anplitude, load-controlled, with sinusoidal axial loading at a frequency of 5 Hz. An extensaneter fran Schenk with a gage length of lOmm was used to measure the strain. All tests were done at roan tenperature.

1.1 Ultrasonic-insitu (USIS) The ultrasonic method (US) is a well established NDI technique to

evaluate the degree of dam3.ge irrparted to catpOsite materials. In this investigation special ercphasis was placed in the use of the ultrasonic pulse-echo mode. For this purpose we developed the USIS method. USIS has many advantages over the ccmron used US-inspection. First, it is possible to do the US inspection of specimens without rezooving them fran the loading frame, i. e. an US inspection of the specimen under load is possible. Second, since a thin fluid film is used as coupling media between US transducer and specimen, an imner­sion in a tank of water is superfluous and avoids water infiltration in the damaged area of the specimen. Fiber reinforced epoxy materials are hygroscopic and the size of the registrated dam3.ged area will change depending on the imnersion time /3/. Finally, the method is time saving and econanic. The US-equipnent consists of a Krautkdimer KB6000 and a selfmade scanning mechanism, Fig. 2. The steering of the scanner is done by an IBM AT, personal carputer, and the management of the data in an IBM 4381 carputer. The results of the US measurement are a standard C-scan and a H-scan or F-scan, developed in the institute. In the H-scan presentation, the distribution of the so called isohyp­si, lines connecting areas of equal back face reflection- or rear echo arrplitude (RE) of the US-signal in the pulse-echo method, are drawn using the Unimap /4/ program. Hereby 100% RE means that the tested part is without dam3.ge. The H-scans are the Irost adequate fonn to present the US results of a catpOsite material part, since they do

523

not only show the location of the damage, but they also show-veq clear-the damage developnent in the loaded specimen. As a matter of fact, the possibility to quantify the damage is given. The F-scan which is similar to the H-scan enables us to present the results of the US examination in a 3 dimensional colored picture.

2 - EXPERIMENTAL RESULTS static tensile test results of speclinens with idealized disconti­

nuities are shown in Tab. 1. An exanple of H-scan pictures at the beginning and after a tensile loading up to 0,6 % longitudinal strain for the stacking sequence 2 as well as the the difference H-scan resulting of both, are shown in Fig. 3. The virgin specimen already reveals differences of the US-RE between above 90 and below 70 %. After loading, the H-scan shows the changes specially in the region of the discontinuity, but also all over the spec.imen. This means, that there are usually potential nuclei of damage distributed throughout the material which can develop to a catastrophic damage depending on stacking sequence and stress conditions. The influence of the stacking sequence upon the damage developnent in spec.imens subjected to static tensile loading, is shown in Fig. 4. The differ­ence H-scan show that the damage developnent in laminate 1 is mre transverse to the awlied load than in the case of the other lamina­tes. Laminate 3 shows the smallest changes over the speclinens. Other investigations made by the authors /5/ with specimens having the same lay-up under tension-tension fatigue loading, show that the nuclei can develop to bands. Many of these bands have a stable growth before a local delamination awrears. These bands are for laminates made out of unidirectional prepregs prilnarily a sign of matrix degradation. Damage developnent in catIX>Site materials are related to stiffness changes, as shown by various authors /6, 7, 8/. The stiffness 0b­tained ·fran the local strain measurement with the extensaneter showed the same tendency as the difference H-scans, i.e. the lay-up with the lowest stiffness showed also the biggest differences in the H-scan.

The stringer stiffened panel was fatigue 2tested under carpression­ccnpression loading (R=0,1 c: =-160 N/mn), sinple suWOrted and without an anti-buCkling devici. Fig. 5 shows the H-cans obtained at N=O and after N=1,5 million cycles. The damaged zone caused by the different inplcts do not behave in the same manner in all cases, dem:lnstrating that the damage developnent is also dependent upon the location of the inplct. The biggest enlargement of the inplct damaged zone was transverse to the awlied carpression load, at the inplct location 3. After about 30000 cycles the stringer, adjacent to inplct 1, begun to buckle elastically due to the degradation of the material at the junction between stringer and panel. Inpact 2 showed no influence concerning the damage developnent in the panel. Stiffness measurements after different nunber of cycles at locations IM, RM, etc., of the panel as function of the corresponding H-scan for the defined region (15x15mn) are also presented in Fig. 5. The measured points show a correlation between stiffness degradation and US-RE diminution.

524

3 - <XJNCLUS IONS The static and fatigue test results presented here show same of

the possibilities of the new US-insitu developed in the institute. The H-scan presented here opens the possibility of damage detection in CFRP before a delamination or catastrophic failure occurs. A feasible way to correlate stiffness degradation and US-RE amplitude was shown.

Aknowlegment The authors express their sincere appreaciation to the members of the computer centre of the DFVLR, Stuttgart, specially Mrs. G. ott.

4 - REFERENCES 1. R.M.Aoki, in AGARD-CP-355, "Characterization, Analysis

and Significance of Defects in Corrposite Materials" (1983) 11 2. D.Wurzel, S.Dehm, in Proc. ICAS 86-4.6.2 (1986), London 3. G.P.Sendeckyj, in AGARD-CP-355, (1983) 2 4. Uniroo.p , Interactive Mapping System, UNIRAS, Lyngby

Derunark 5. R.M.Aoki, J.Heyduck, DFVLR-IB 435-88/18 (1988) 6. S.W.Tsai, H.T. Hahn, Corrposite Mat. Workbook, AFML-TR-77-33 7. R.D.Jamison, et.al., ASTM STP 836 (1982) 21-55 8. T.K.O'Brian, ASTM STP 876 (1985) 282-297 9. K.L.Reifsnider, W.W.Stinchcomb, ASTM STP 907 (1986)

298-313

No.

1 2 3 4

Laminate

(0/±45/90) (±45/0/90)s (0/90/±45)s (90/±45/0):

Table 1

trnax. (%)

0,726 0,64 0,76 0,72

T30

0/C

OD

E69

G

EW

.

[ 021

± 4

5/90

/ ± 4

5/0z

1 s

X

I I I I

-........

I A

O

0 I ®

LM

I I I

I I I I

RM

8:

LU

I I

Fig

. 1

Str

ing

er

sti

ffen

ed

pan

el

wit

h

imp

act

dam

aged

zo

nes

(1

,2,3

)

m

'ft

'. iI

*

4*t~

t~ ~

.- Fig

.2 U

S-i

nsi

tu:

Scan

ner

WA

DIN

G F

RA

ME

i.

,....

TRA

NS

OO

CE

R

SP

Er:

IME

N

(Jl

N

(Jl

526

30~----~~~--~~~~~=--1~--4

25 +-----+---~~---+----~----4---~ E E 20+-----+---~~~

~ 5~----~----TP--~~----_+----_+----_1

10 +-----~~~~----~~--r_----r_--~

5+---~~~~----_+----~----4_--__1

O+-~~~SL~~--~r_----~~~r_--~

35-W~:!I;t1!II

30

25 E E20-fI:.l.~~

~ 15-+---"--'r.!!!I'

IO--t;,1'if'f'j~

o 10

10

20

20

x mm

30 40 50 bO

X mm

30 40 50 bO X mm

Fig.3 H - Scans

(±45/0/90)s

E = 0 %

H-Scan after

loading up to

£ = 0.6 %

CJ AeovE 90.0

CJ 8~,0 - 90,0

eo.o - 85,0 - 75.0 - 80.0 - 70.0 - 75,0 -8aow 70.0

Difference H-scan

f = 0% and f =0 .6%

c:::: ABOVE 20.0 CJ ~,o - 20.0 o 10,0 - 15.0

5.0 - 10.0 _ saow 5.0

35

30

10

5t:;i.~~ o-""'"-~~ ..

25 E E 20

>-- 15

10

5

10

x mm

x mm

20 30

X mm 40 50 60

527

Fig. 4

Difference H-Scans

(0/±45/90)5

E=O% and £=0.71%

(0 / 90/±45) 5

€=O% and (=0.71 %

~90/±45/0)s

£=0% and £=0. 70%

o ABove 20,0 o 15,0 - 20,0

D 10,0 - 15.0 _ ~,O - 10,0 _ saow ~.O

528

1>0

50

40 E E >-.,30

20

10

.DO .. 10 .. ?O

10 UJ a: • OIl

til ::::> ..,

3D

2D

••

•

2D

x mm

x mm

00

0

u

0

0

0 0

0

STIFFNESS GPa

c

.. .00

Fig. 5 H - Scan

Stringer stiffened panel

with impact damages

N = 0

H - Scan after

N

'-I --o -

1.5 106 cycles

.... BOvE QO 70 · 90 so · 10

'" 50 10 )0

ea.Ow 1()

~sured stiffness and

H - Scan evaluation

STIFFNESS CHANGES DURING FATIGUE OF ANGLE-PLY GLASS/POLYESTER OF HIGH QUALITY UNDER

VERY LARGE NUMBER OF CYCLES

S. ANDERSEN, H. LlLHOL T

Riso National Laboratory Metallurgy Department - 4000 ROSKILDE - Denmark

ABSTRACT

Glass/polyester materials of high quality have been studied under long time fatigue loading, with special reference to the associated stiffness and strength changes. The continuous recording of stiffness parameters during fatigue testing constitutes a non-destructive method by which microstructural changes can be monitored. The combined infor­mation on the material modulus E and the secant modulus Es allows a schematic stress-strain curve to be invoked.

INTRODUCTION

Glass fibre reinforced polyester is a material with generally useful properties, which can be achieved at a reasonable price. In particular, the deliberate placing of fibres along the heavy loading directions will give a material with good properties on a weight basis, which can compete with and even supersede conventional metals like steel and aluminium alloys. A good example is the extensive use of glass/polyester as a material for large wingblades for rotors on windturbines. Under the conditions of wind and gravity the loads are high and the planned and expected lifetime is long. The particular load history is fatigue due to irregular wind oscillations and (re­gular) variations of the gravity force during rotation of the rotor. The lif~ time is expected to be about 20 years, corresponding to more than 10 cycles.

Under these long time fatigue conditions the glass/polyester material is sensitive to defects, both originating from the fabri­cation and from the service of the windturbine. These defects, typi­cally cracks, reduce the stiffness and strength, and thus the design

530

basis for the lifetime of the wingblades.

A series of fatigue data for failure of the material are shown in fig. I, which illustrates the importance of the fibre orientation and of the material quality.

The fibre orientation for angle-ply laminates beyond about 10· reduces the fatigue strength and fatigue limit significantly. The material of ±5· orientation is of lower quality, in particular measured as a higher porosity content, and the fatigue curve is placed below the 10· curve.

The present study is aimed at measuring and understanding the development of cracks during fatigue loading and the associated change (reduction) in stiffness and change of the stress-strain curve of the material.

I - MATERIALS

The glass fibre reinforced polyester has a fibre content of 50 vol% and an (initial) porosity of less than 0.5%. The fibre orien­tations, selected for detailed study, are angle-ply of flO· and ±60·.

II - FATIGUE TESTING

Tests were performed in tension-tension fatigue, with R=O.l, at a frequency of about 5 Hz. The tests were load controlled and taken to final failure of the specimens (which is not a well-defined state of final damage). During the fatigue testing the material's E-modulus was measured by an (in-situ) tensile test to low strain. The secant modu­lus Es was recorded as the stress-range divided by the instantaneous strain-range. Both E and Es were recorded at every tenth of a decade during testing, and plotted as a function of number of cycles.

The stiffness data are plotted in a normalised diagram, with E/Eo and Es/Eo respectively as a function of log N/log Nmax ' where Eo is the initial modulus at the start of testing and Nmax is the number of cycles at failure. These types of plots allow an easier and clearer comparison.

III - MATERIAL MODULUS E

For the two glass/polyester materials of flO· and ±60· fibre orientation the values of E/Eo are shown in fig. 2 and 3, each at three different levels of fatigue loading, measured as the maximum (initial) strain of the load cycle. The normalised plot allows a di­rect comparison of the three strain levels, even if the corresponding lifetimes are very different.

The strain levels for the two materials are rather different, but the fatigue curves of fig. 1 can be used to illustrate a correspon-

531

dence in terms of lifetimes, as shown in table 1.

Table 1 Material

strain level ±10· ±60· log Nm,x (fig. 1) "high" 1.18% 0.32% - .2 "medium" 0.93% 0.20% -4.6 "low" 0.58% 0.10% >7

The shapes of the curves of E/Eo are very similar for the two materials at the same strain level, as defined in table 1.

A possible interpretation of the shapes of the E/E -curves at the three strain levels corresponding to short, medium and Yong (infinite) lifetimes, is the following:

at high strains, rather few cracks can develop and grow before a crack configuration is formed which results in (final) failure. The lifetime is short and the stiffness reduction is moderate.

at medium strains, many cracks can form and grow before a criti­cal configuration is formed. The lifetime is relatively long and the stiffness reduction is significant, with E decreasing to a level of 70-75% of Eo'

at low strains, which are close to or perhaps below a fatigue limit, relatively few cracks will form, and their growth is re­stricted. The lifetime is very long (infinite) and the stiffness reduction is small and stabilizes at a certain level towards the end of the life. At present it is not clear why the ±60· laminate shows an increase in stiffness before stabilization.

v - MATERIAL MODULUS E AND SECANT MODULUS Es

To study the general shape of the stress-strain curve of glass/ polyester after (same) fatigue loading a detailed comparison between E and Es has been made for the ±10· laminate at medium strain level. Two strain levels of 0.8% and 1.0% respectively are used for the nor­malised plots of E (fig. 4) and Es (fig. 5).

At this medium strain level the reduction in stiffness is gene­rally large, and a comparison between E and Es shows a tendency, which is enhanced schematically in fig. 6. Three regions can be identified and their characteristics are listed in table 2.

Table 2 Region

a b c

Stiffness Lifetime log N/log Nmax o -0.4 0.4-0.7 0.7-1.0

The mechanisms of damage which can be responsible for the ob-

532

served changes in stiffness parameters E and Es will be based on the following types of cracks and their implied effect on the deformation behaviour of the composite material:

matrix cracks parallel to the loading direction and fibre di­rection (±lO·) will not lead to (measurable) changes in E, but may cause some increase in the significance of the matrix visco­elastic behaviour on the (overall) deformation of the composite, i.e. Es<E.

debonding (nearly) parallel to the loading direction and at the interface between fibres and matrix can have two effects (fig. 7): o debonding will most likely start at the fibre ends, causing

the effective fibre length for load transfer to decrease; this will result in a (small) reduction in E.

o the significance of the matrix visco-elastic behaviour on the composite deformation will increase, i.e. Es<E.

matrix cracks at right angles to the loading direction and fibre direction (±lO·) will lead to a reduction in E, but not to any change (increase) in the matrix visco-elastic contribution, i.e. Es=E «Eo)' There may be a contribution from the (matrix) crack opening to the composite deformation, and this will in principle cause that Es<E.

On the basis of these possible types of cracks in the glass/poly­ester, the following interpretation of the observed changes in E and Es can be given:

Table 3 Region

a b c

Crack types (fig. 11) none

3 1 and 2

A possible identification of likely mechanisms of crack formation during the (normalised) fatigue life can therefore be the following, based on the glass/polyester ±lO· laminate.

a: at short lifetimes, log N/log Nmax < 0.4, very few, if any, cracks, have formed, and no (measurable) changes in stiffness can be recorded.

b: at medium lifetimes, log N/log Nmax between 0.4 and 0.7, cracks at right angles to the loading direction develop, and they lead to reduction in E, but to no significant change in the matrix visco-elastic contribution, so that Es=E.

c: at long lifetimes, log N/log Nmax between 0.7 and 1.0, additional cracking occurs parallel to the loading direction (matrix cracks and/or debonding), so that the matrix visco-elastic contribution becomes more significant, and thus Es<E.

The (schematic) stress-strain curves corresponding to this inter­pretation of cracks and deformation mechanisms are shown in fig. 8.

533

This may be a simplified picture. because interactions can occur be­tween the different types of cracks. both their formation. number and growth.

VI - CONCLUSION

The properties of composite materials with a relatively brittle matrix. like glass/polyester laminates. are governed by the formation of cracks. in particular during fatigue. The (continuous) measureme~t

of stiffness parameters. like E and Es. is a convenient (non-destruc­tive) method to record changes in the microstructure of the composite. The combined information from E and Es allows a schematic stress­strain curve to be invoked and thus gives useful knowledge on the strength properties.

------ -------

" ----r-------1----

o ~, - ------------ --

Ll--- 1--l

1. 1---1--1--- -"'- . - •• - .~ ~"---< . , . '" . ~ ----------I~ -- !-

. ~_II_~_~._~_~ __ L_~ __ ~ •

~,.(N)

Fig. 1 . Fatigue data for tension-tension loading at R=O.l. for glass/ polyester with 50 vol% fibres and fibre orientations: o 0'; "J ca.± 5'; # flO'; 0 ±60'

."

, - ~'" '\f-.-... ~ .n --- - - - ------ ~S'l'! ;;,

,1:~ --------- - -----

, •

--- -----'---,----' . ' .. 'lJg(N)/lgv(H"mx)

Fig. 2. Material stiffness E as a function of N. in normalised diagram. for glass/polyester with flO' fibre orientation; maximum initial strain: o l.18% ; X 0 . 93%; - 0.58%

... 1---1-------

~---

.~ __ L_ __ L__~

tI .2: 11 .8

log(N)/ lrry(H"uu:)

Fig. 3. Material stiffness E as a function of N. in normalised diagram. for glass/polyester with ±60' fibre orientation; maximum initial strain: o 0.32% ; X -0.20%; - 0.10%

534

o

''''. fL~!:/J1()I,yp. .~t(l''- I (J-fl,' ,I./"" " IIfJlllinnl , ' rn i" I,.",.t ." ("': 1' ----- - - .-. ...

~ ... , ---f--._- - _ . ---.. -----

0'----'----·,- -- --- _ .. ---. LGJ/(II)/I,vg(Nllut.Z)

Fig. 4. Material modulus E as a function of N, in normalised dia­gram, for glass/polyester with ±10' fibre orientation; medium strain levels of 0.8% and 1.0%.

h

n . os 1 . 0

Fig. 6. Schematic curve for E.and Es (related to fig. 4 and 5) with three regions a, band c, as de ­scribed in the text.

Fig. 8. Schematic stress-strain curves for glass/polyester, after fatigue loading corresponding to regions a, band c, as described in the test and fig . 6 .

. -.

(;l,u.!,=.o,:-/ ,Jul!ll '.-.:,pr 10 -· ,/ r·y'-'·(· N'''lIu.u' .~ r. nl" ' .. t·"'. I"':)

Fig . 5. Secant modulus Es as a function of N, in a normalised diagram, for glass/polyester with ±10' fibre orientation; medium strain levels of 0.8% and 1.0%.

L __ , ~ c=---'

CC-J) Fig. 7. Schematic microstructure with cracks: parallel cracks (1), debonding (2), and cracks at right angle (3) to the loading direction.

a

c

3D-FABRICS FOR COMPOSITE SANDWICH STRUCTURES

I. VERPOEST, M. WEYERS, P. de MEESTER

Katholieke Universiteit Leuven Department of Metallurgy & Materials - De Croylaan 2 - 3030 LEUVEN - Belgium

ABSTRACT

The skin-to-core adhesion strength in sandwich panels is very often the weak link in composite structures. For this problem, a novel solution has been developed, based on a new type of fabric. A three-dimensional fabric is woven in one step; pile threads connect the top and bottom twodimensional fabric. When the 3D-fabric is impregnated, cured and filled with a PU-foam, a sandwich structure with interconnected skins is created. Preliminary mechanical test results are encourageing.

Keywords: composite materials, fabrics, , sandwich structures.

INTRODUCTION

Composite laminates are the favorite materials of aeronautical designers: under in-plane loading conditions, they show an extremely high specific strength and stiffness, so that important weigth reductions can be realised in airplanes and space structures. Where bending stresses are predominant, a further weigth reduction is realised by the use of sandwich panels: the load carrying composite laminates are separated by an weaker, but extremely light core, so that the structural bending stiffness is increased. This core can be either a foam or a honeycomb; honeycombs , made out of thin nome x paper or aluminum foils, are preferred mainly for their better shear properties, but are more expensive and more difficult to machine than foams.

536

Out-of-plane stresses can however create serious problems in sandwich structures. In honeycomb as well as in foam core sandwich structures, the thin and stiff skin is adhesively bonded to the core. Resistance against localised compressive forces is low. Impact of a foreign object on the skin can hence introduce damage to the core, and debonding between skin and core. Delaminations can also develop from a stress raiser. Both these delamination types can spread out easily during subsequent bending; the delaminated skin will buckle away from the core when the skin is at the compression side. Improving the delamination resistance between skin and core is badly needed for this type of structural elements.

The weak skin-to-core bonding in sandwich structures still is an important weak link in composite sandwich structures. It is a similar problem to the low value of the interlaminar strength in composite laminates.

In this paper, a solution will be presented for both these problems. This novel solution is the result of a close cooperation between a weaving company, looking for new application fields for their high technology weaving techniques, and some researchers of the Composite Materials Group K.U.Leuven. This paper reports on the exploratory studies which have been carried out during the past two years.

It will be shown how a simple three dimensional fabric has a potential to improve the delamination resistance of sandwich structures. In a future paper {l} , it will be demonstrated how a two - and - a - h a I f dimensional fabric, made from the former 3D-fabric, enhances the delamination resistance of laminates.

MANUFACTURING 2,5- AND 3D-FABRICS

The 3D-fabrics consist of two twodimensional fabrics, which are linked to each other by pile threads. The major difference with stitched composites is that these 3D-fabrics are woven in one step (fig 1): the weft threads ( or "rovings") for the upper and lower fabric ( or "backing") are brougth separately into the weaving loom; the pile threads are connected to the warp threads, and so to the whole backing. The pile density can be changed by altering either the warp density or the "connection frequency" with the warp threads. Moreover, the pile heigth can vary continuously between 4 and 25 mm. The longitudinal stiffness of the fabric uniquely depends on the density and material of the weft threads.

{To make 2,5D-fabrics, the 3D-fabric is simply cut into an upper and a lower part, immediately after leaving the loom. -The name 2,5D-/abrics is chosen, because after cutting the pile threads only point out into half of the third dimension; one could also call it a "hairy" fabric ! -}

For this exploratory research project, basically two types of 3D-fabrics have been manufactured . - Their weaving patterns, and the materials used, are described in Table 1.- The top and bottom backings are woven with 2x68 tex E-glass rovings in weft and warp; the pile is aramid (Twaron HM from AKZO Co.) or E-glass. The width of the fabrics varies between 45 and 65 mm.

537

MANUFACTURING 3D-FABRIC SANDWICH STRUCTURES

First, the 3D-fabric is impregnated with an epoxy resin ( Shell Epikote 816 + Epicure DX161, mixing ratio 100/50). Then, the 3D-fabric ( length = ±. 500mm, width 50 to 60 mm) is mounted in a frame, and a small tension load is applied in the longitudinal direction, in order to stretch the fabric. In this frame, the 3D-fabric is open: the pile fibres are almost vertical in between the top and the bottom fabric.

The frame is brought into an oven, and the impregnated, open 3D-fabric is cured at 100°C for 60 minutes. After cooling, a polyurethane foam is injected into the 3D-fabric. It immediately expands, filling the free space between the pile fibres. So, a sandwich structure is formed with similar density to the usual Nomex honeycomb sandwich structures ( ±. 0,15 kg/dm3). More detailed information on the manufacturing procedure can be found in {2}.

MECHANICAL TESTING

The faces of sandwich structures are thin, and hence out-of-plane loads will have to be caried partly by the core. So, compression tests ( according to ASTM C 365) have been performed on the pure foam, on the empty and on the foamed 3D-fabric (specimens 3D-I and 3D-2, Table I). The load was applied perpendicular to the plane of the top and bottom fabric; square specimens of 60 by 60 mm were used. As most sandwich structures are loaded in bending, four point bending tests were carried out as well ( according to ASTM C 393). Finally, in order to evaluate the well known weakness in shear of foam core sandwich structures, shear tests have been carried out ( according to ASTM C 273).

For comparison reasons, the aforementioned tests were also performed on sandwich structures containing a Nomex honeycomb (heigth 12,6 mm, hexagonal 1/8" cells). The skins were lrecured single 2D-fabric layers; for adhesion, a structural film Structufilm R-382 was used (both honeycomb and film are Hexcel Co. products).

RESULTS AND DISCUSSION

DENSITY In order to be competitive with normal sandwich structures,

the density of the 3D-composites should not exceed the honeycomb sandwich density. Fig. 2 shows that only the 3D-composites with high density glass fibre piles have a slightly higher density. Moreover, this figure indicates that a low density foam was used in the 3D-fabric composites.

COMPRESSION As the pile fibre bundles are impregnated with epoxy resin,

538

they have already a certain compression strength on their own. However, they are very slim, and will buckle easily. After foaming, the polyurethane foam will sidewise support the pile fibre bundles. Hence, buckling is prevented and the piles can have a reinforcing effect on the foam. Fig. 3 clearly shows that adding pile fibres to the foam increases the compression strength with a factor of 3 (low density aramid pile) to 5 (high density glass pile) !

The maximum attained value however is still less than half of the compression strength of honeycomb sandwich structures. The main reason for this difference is the weakness of the foam used in this test series. In future tests, superior PU-foams (with compression strengths up to 0.6 MPa) will be used, so that the foam reinforcement by the piles should be more effective as well.

FOUR POINT BENDING The flexural stiffness EI has been calculated from the four

point bending test results. Fig.4 shows that the 3D-fabric composites are clearly superior to the honeycomb sandwich structure. The main reason for this improvement is that in the 3D-fabric composites the pile fibres also partly reinforce the skin fabrics. As the pile density increases, and hence the fibre volume fraction in the skin, one can observe that also the flexural stiffness increases substantially. These pile fibres are completely absent in the top and bottom skin of the honeycomb sandwich structures.

SHEAR TESTING Preliminary results of the shear tests indicate (fig.5) that the

shear strength and stiffness of the 3D-fabric composite is 3 resp. 5 times lower than the corresponding values for honeycomb sandwich structures. This result reflects the intrinsic weakness in shear of all foams used as sandwich cores. Further experiments will be carried out to check the effect of stronger and stiffer foams.

In contrary, the shear energies (up to total failure) are almost equal; the reason for this being the fact that the pile fibres hold the two parts of the 3D-fabric composites together, even after shear failure of the foam occurred.

CONCLUSIONS

A new type of fabrics has been presented: they can be used as 2,SD-fabrics, and as 3D-fabrics. It was demonstrated that sandwich structures with intimately bonded skins can be produced from these 3D-fabrics. Some mechanical properties do not yet equal those of honeycomb structures: the compression and shear strength are too low, but it is believed that this problem can be overcome by using a better foam. Flexural stiffness and density however are equal to or better then honeycomb sandwich structures. More research is needed to further prove the other expected advantages of this new material: high impact resistance, high bolted joint strength and improved delamination resistance of the sandwich structures based on these 3D-fabrics.

539

ACKNOWLEDGEMENTS

The staff members of the Schlegel Co., mr. Demarest , Mr.Boey and Mr. Declercq, are kindly acknowledged for their coOperation, and for providing the fabrics. Most of the experimental work was carried out by students during their engineering thesis; we thank Yves Bonte and Johan Swaelen for their essential contribution.

REFERENCES

1. I. Verpoest et.al., 2.5D- and 3D-fabrics for delamination resistant composite structures, to be presented at ICCM-VII, Bejing (China), august 1989 2. J.Swaelen, Optimisation of 3D-fabrics for composite sandwich structures, Ind. Engineering Thesis, KIH.De Nayer, Mechelen (Belgium), 1988 (in Dutch)

TABLE 1. FABRIC DESCRIPTION

-------FABRIC WARP+WEFT PILE PILE PILE

<IIE MATERIAL MATERIAL LENGTH (mm) DENSITY

(cm-2)

3D-l E-glass E-glass 20 18

2*68 tex 2*68 tex

3D-2 E-glass Aramid HM 20 2,5

2*68 tex 1260 dtex

TOI'WEI'T

1I0TTOMWEFT

Figure 1: weaving scheme for 2,5D- and 3D-fabrics (longitudinal cross section)

540

0.25

0.2

0.15

0 . 1

0.05

o

DENSITY (g/cm3)

GLASS- ARAMID-HIGH DENS LOW DENS

HONEY­COMB

PU -FOAM

Figure 2: Density of the 3D-fabric composite, the honeycomb sandwich and the foam.

2

o

COMPRESSION STRENGTH ( MPa )

1.9

GLASS­HIGH DENS

0.17

ARAMID - HONEYCOMB PU-FOAM LOWDE S

Figure 3: Compression strength of the 3D-fabric composite, the honeycomb sandwich and the foam.

FLEXURAL STIFFNESS EI ( m2)

154

o

Figure 4: Flexural stiffness of the 3D-fabric composite and the honeycomb sandwich.

5

4

3

2

o

3D-FABRIC mIil HONEYCOM B

SHEAR STIFFNESS

(GPa)

SHEAR STRENGTH

(MPa)

SHEAR ENERGY ( IOJ )

541

Figure 5: Shear properties of the 3D-fabric composite and the honeycomb sandwich.

FLUAGE CREEP

Chairman: Dr A. R. BUNSELL Ecole Nationale Superieure des Mines de Paris

NON LINEAR VISCOELASTICITY APPLIED FOR THE STUDY OF DURABILITY OF POLYMER MATRIX COMPOSITES

ABSTRACT

A. CARDON, H.F. BRINSON', C.C. HIEL"

Free University Brussels - Pleinlaan 2 - 1050 BRUSSELS - Belgium 'University of Texas at SAN ANTONIO TX 78285-0665 - USA

'WASA - Ames Research Center - MOFFET FIELD CA 94035 - USA

For structural applications an essential problem is the analysis of the long term behaviour under a general, mechanical and environmen­tal, loading history. This durability analysis, or life-time prediction, must be performed on the basis of a certain number of tests carried out on limited, and if possible short, time scale by the use of accelera­ting factors. Acceleration of the thermomechanica1 behaviour of polymer based systems can be realised by temperature and stress level. In order to analyse the results of such tests nonlinear viscoelasticity theory must be used.

INTRODUCTION

The viscoelastic nature of the polymer matrix combined with the elastic nature of the fibers and the specific properties of the inter­face region fibers-matrix, induces an anisotropic elastic-viscoelastic behaviour in the unidirectional reinforced composite. Under quasi sta­tic loading conditions the composite exhibits time effects and under cyclic loading frequency effects. The mechanical behaviour of the composite is not only function of the loading, and consequent stress distribution, but also function of tem­perature, moisture and other environmental conditions. If the visco­elastic nature of the composite appears especially by loadings perpen­dicular to the fiber direction and off-axes shear loadings, a composite laminate is ~t any moment subject to stress transfers from matrix to fibers and from one layer to another.

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I - DURABILITY ANALYSIS

1.1. Basis

Some elements of a general loading can be used in order to acce­lerate the behaviour of a viscoelastic continuum. Such accelerating factors must perform some equivalence between high levels during short times and lower levels for long times. For simple thermorheological media, applicable to many polymers, such an "equivalence" exist between temperature and time : the well-known time-temperature superposition principle, (TTSP ; WLF ; master curve). H. Brinson and co-workers, /1,2,3/, showed the applicability of the TTSP to unidirectional rein­forced polymer matrix composites. A second possible accelerating factor is the stress level. This possi­bility was experimented by the research group of H. Brinson at Virginia Tech from 1981 on. Collaborative efforts in the same domain were made at the composite research group of the V.U.B. and especially by C.Hiel, /4/, and R. Brouwer, /5/. The possibility to use high stress levels during short periods for prediction of the behaviour at lower stress levels for long terms is the basis of the Time-Stress-Superposition-Principle (TSSP) and in com­bination with temperature the general Time-Temperature-Stress-Superpo­sition Principle (TTSSP).

1.2. Analysis of the short term results - nonlinear viscoelasticity

The viscoelastic behaviour under transverse and shear loading of a unidirectional reinforced polymer matrix composite layer at higher temperature and stress levels becomes rapidly nonlinear. This implies the need of nonlinear constitutive equations in order to analyse the experimental short term results. Many nonlinear viscoelastic equations exist but most of them, and es­pecially all the multiple integral formulations are very difficult to use for the analysis of experimental data. Single integral formulations are more convenient and the thermodynamic based theory proposed by R. Schapery, /6/, is a good framework for this analysis. Starting with the construction of the master curve from tests at dif­ferent temperature levels within a limjted'frequency range, combined with isothermal creep and creep-recovery test results, the complete nonlinear description of the viscoelastic behaviour of the composite is obtained. The prediction of the evolution of the mechanical cha­racteristics of a unidirectional reinforced lamina is obtained and by lamination theory the durability analysis of a laminate is possible.

1.3. Schapery's nonlinear viscoelastic model From the general expressions

aGR, aF flji na = - ar- + ~ 6Sqr

a a-oo

547

with na generalised coordinates

Fa gene ra 1 i sed fo rce s GR Gibbs free energy in the reference state

Fq internal functions of Fa

dt dt dljl = -- = - reduced time a a (~) 0 aG

we obtain the one-dimensional formulation :

fljl d [g2°] E = go So 0 + gl 6S(ljI-ljI') ~ dljl'

J -00

where the nonlinearising functions go' gl' g2 and ao may be functions of temperature, stress and other factors. The entropy is controling ao by aD· The complete analysis needs the measurement of the variations of the nonlinearizing functions and therefore creep and creep-recovery data are necessary.

1.4. Applications

This methodology for a durability analysis was applied to thermo­set matrix composites under uniaxial and biaxial loadings, /3/,/4/,/5/. The obtained results were satisfactory on the basis of comparison with the real time behaviour on a limited time range, /3/. For thermoplastic matrix composites the analysis was developed by X.Xiao, /7/,/8/. The prediction of long term behaviour is still pos­sible even if there where some difficulties with the different nature of the creep-recovery in transverse and in shear conditions. This spe­cial behaviour may come from the anisotropic nature of the interface region fiber matrix in APC-2.

2. EXPERIMENTAL RESULTS

We limit here the presentation of results to some of them obtained on APC-2. Fig. 1 gives the compliance of APC-2 at different temperature levels for a limited time or frequency range: (a) for transverse ten­sile and (b) in shear, from off-axes testing. Fig. 2 gives the evolution of the shear compliance by application of the TTSP for APC-2. Fig. 3 gives the final results, after use of the isothermal creep and creep recovery data, of the evolution of the shear compliance of APC-2 by application of TTSSP.

548

transverse tensile APC-2

'" u c:

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6 u L

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Fi gure 1

Lo~ time (min)

Fi gure 2

shear APC-2

APC-2 (TT)

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o,,~ <.') ...... -~ u c

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CONCLUSIONS

TTSSP is applicable, with the use of Schapery's nonlinear visco­elastic model for the analysis of the short term test results in order to obtain the evolution of the mechanical characteristics of a unidi­rectional lamina and gives us the basis of a durability analysis of a general composite laminate under different mechanical and environmental loading histories.

In order to have an analysis valid for a general loading, inter­action with fatigue behaviour must be studied and damage evolution must be measured.

REFERENCES

1 - Brinson H.F., "Experimental mechanics applied to the accelera­ted characterization of polymer based composites". New Trends in Expe­rimental Mechanics (Ed. J.T. Pindera), 1981, Springer Verlag, Vienna, pp. 1-69.

2 - Brinson H.F., "Viscoelastic behaviour and lifetime, (durabili­ty), predictions". Mechanical Characterisation of Load Bearing Fibre com~osite Laminates (Ed. A.H. Cardon - G. Verchery), 1985, Elsevier App ied Science Publishers, London - New York, pp. 3-20.

3 - Tuttle M.E. and Brinson H.F., "Prediction of the long-term creep compliance of general composite laminates", Experimental Mechanics, r~a rch 1986, pp. 89-102.

550

4 - Hiel C.C., "The nonlinear viscoelastic response of resin matrix composite materials". Doctoral thesis, V.U.B., 1983.

5 - Brouwer R., "Nonl inear viscoelastic characterization of trans­versely isotropic fibrous composites under biaxial loading". Doctoral thesis, V.U.B., 1986.

6 - Schapery R.A., "On a thermodynamic constitutive theory and its application to various nonlinear materials", Thermoinelasticity (ed. B.A. Boley), Springer Verlag, Wien - New York, 1970, pp. 259-285.

7 - Xiao X., "Viscoelastic characterization of thermoplastic matrix composites (APC-2)", Doctoral thesis, V.U.B., 1987.

8 - Xiao X., "Studies on the viscoelastic behaviour of a thermo­plastic resin composite" (accepted for publication in Composite Science and Technology).

COMPORTEMENT AU FLUAGE DE STRATIFIES POLYESTER VERRE E DESTINES A DES APPLICATIONS NA V ALES

A. LAGRANGE, R. JACQUEMEr

IFREMER - BP 70 - 29263 PLOUZANE - France 'ATOCHEM - Cerdato S3T - 27470 SERQUlGNY - France

Comportement au fluage de stratifies polyester Verre E destines a des applications navales

R. JACQUEMET (ATOCHEM) A. LAGRANGE (IFREMER)

ABSTRACT

A comparative study of durability in distilled water of two various polyester laminates reinforced by fibre glass type E is presented. A three point bending apparatus has been developped allowing immersion of samples in distilled water at various temperature and at high stress level. (30, 50 and 70 % of the nominal breaking point). Results show that tetrapolyester laminate have a least resistance (failure after quite short time and elevated rate of deflection) than isopolyester laminate. At same level stress, time rupture of samples and deterioration mechanisms depend of the orientation of the fibre. The prediction of the long time comportment seems more exactly by use an Arrhenius's model than by extrapolation of tests realized at high stress level.

552

INTRODUCTION

Les materiaux composites a matrices polymeriques et a fibres minerales ou organiques sont de plus en plus utilises dans des domaines divers comme l'aeronautique, l'offshore petrolier ou la construction navale.

En technologie navale, la nouveaute du materiau (absence de donnees sur son comportement a long terme) et la meconnaissance des char gements reels appliques sur une structure marine, amement les concepteurs a utiliser les materiaux composites avec des coefficients de securite eleves [1J . Dans ce secteur industriel, les resines polyester sont couramment employees pour la fabrication de coques de bateaux et comme protection de celles-ci par l' intermediaire d 'un gelcoat. Actuellement la prevision du comportement a long terme sur ce type de resine se heurte a une absence de donnees contrairement au cas des resines epoxydes qui ont fait l' objet de nombreux travaux ces dernieres annees (2J [3J

La prevision du comportement a long terme passe par l'etablissement de protocoles d'essais de vieillissement acceleres representatifs des conditions reelles d'utilisation. lIs permettent de mettre en evidence, de qualifier, de quantifier les degradations physiques ou chimiques, (plastification, gonflement, hydrolyse) reversibles ou irreversible6et d'effectuer des previsions de duree de vie. La finalite de cette etude consiste donc a comparer Ie comportement au vieillissement sous charge de 2 stratifies polyester isophtalique et tetrahydrophtalique en immersion a hautes temperatures.

I MATERIAUX ETUDIES

1.1 - RESINES

2 types de resines (NORSOLOR) ont ete choisies

1 resine polyester tetrahydrophtalique (TETRA) 1 resine polyester isophtalique (ISO)

1.2 - RENFORT

Les structures navales sont constituees de materiaux composites de grandes diffusio~ avec un renfort de fibre de verre E (tissu equilibre) afin de limiter Ie coOt de la matiere premiere.

Afin de faciliter la comprehension et la modelisation des differents processus de vieillissement, Ie renfort de fibre de verre E choisi est a armature quasi unidirectionnelle (90 % sens chaine et 10 % sens trame) de 230 g/m2 a ens image textilo­plastique. Le taux de renfort est de 50 % en masse.

553

1.3 - MISE EN OEUVRE

Les coques de bateaux sont fabriquees par la technique de moulage au contact a froid. Par analogie avec ce procede de mise en oeuvre , les eprouvettes sont elaborees au eETIM de Nantes par moulage au contact a froid (20Q e), sous presse, sous forme de plaques de 400 x 400 x 5 mm3.

Une post-cuisson de 2 heures a 100Q e, optimisee par analyse enthalpique differentielle, permettant de stabiliser et de reticuler les resines a 80-90 %, est effectuee sur chacun des materiaux.

II ESSAIS DE FLUAGE EN FLEXION 3 PANNES

II.1 - DESeRIPTIF EXPERIMENTAL

Des bancs de flu age ont ete con<;us pour appliquer une charge constante sur l'eprouvette, sollicitee en flexion 3 pannes et pouvant etre exposee a l'air ambiant ou en immersion a differentes temperatures (figure 1)

Figure 1 Schema de principe du banc de fluage en flexion 3 pannes permettant 5 essais simultanes.

L'eprouvette qui repose sur deux appuis reglables en longueur est soumise a une charge, par l'intermediaire d' un arbre guide en translation verticale et situe a egale distance des deux supports inferieurs.

554

Un capteur de deplacement situe a l'extremite superleure de l'arbre permet de suivre la fleche au centre de l'eprouvette en fonction d4 temps par l'intermediaire d'une acquisition automatique de donnees sur une centrale SCHLUMBERGER de type ORION.

11.2 - DIMENSIONS DES EPROUVETTES

Les dimensions des eprouvettes sont definies ci-dessous

sens longitudinal :

sens transversal

largeur : 30 mm epaisseur : 5 mm distance entre appuis largeur : 30 mm epaisseur : 5 mm distance entre appuis

II.3 - MILIEUX AMBIANTS ET PARAMETRES MESURES

145 mm

95 mm

Les bacs sont alimentes en eau distillee et les eprouvettes sont immergees a differentes temperatures (40, 60, 8011 C) afin d'activer thermiquement les differents processus de degradations. Les temperatures d' essais sont bien inferieures aux temperatures de transition vitreuse, de telles facons qu'elles n'induisent aucune degration thermique. Les experiences sont realisees a differents taux de sollicitation (30%, 50% et 70% par rapport a la contrainte a la rupture) dans Ie sens longitudinal et transversal du renfort de fibre de verre apres determination des contraintes a la rupture par essais mecaniques statiques.

Deux parametres sont mesures :

- La vitesse de fluage € (en mm/heure) correspondant a la pente de la zone lineaire A - B de la courbe de deformation en fonction du temps (figure 3a)

- La duree de vie tr (temps a la rupture) du materiau.

III RESULTATS ET INTERPRETATIONS

L'influence d' une charge mecanique sur un materiau cree un endommagement qui augmente la quantite d'eau absorbee par Ie materiau a saturation (figure 2) (4::1 Le taux de saturation dans un materiau composite depend du taux de contrainte, du sens de la sollicitation par rapport au renfort, de la nature chimique de la resine et de la temperature du milieu d'absorption. Les tests de fluage en flexion 3 pannes, realises sur ces deux stratifies, en immersion a dHferentes temperatures (40, 60, 8011 C) mettent en evidence Ie meilleur comportement du stratifie ISO en immersion et sous contrainte vis a vis du stratifie TETRA. La duree de vie du materiau est plus €!levee pour Ie strati fie ISO, alors que sa vitesse de fluage est plus faible, vis a vis du stratifie TETRA (figures 3a et 3b). Enfin l'accroissement du taux de contrainte et de la temperature du milieu augmente la vitesse de fluage et redui~ la duree de vie du materiau. Figure (4a et 4b)

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Figure 2 : Influence des charges mecaniques en flexion 3 pannes sur la cinetique de sorption du polyester isophtalique!verre E en eau distillee a 40 QC.

Pour un taux de contrainte de 50 %~RF, la duree de vie du stratifie TETRA est plus courte dans Ie sens longitudinal que dans Ie sens transversal du renfort a 60Q C alors qu'elle est superieure dans les sens longitudinal a 40Q C (tableau 1). Cette inversion de la duree de vie et de la vitesse de fluage entre les deux sens de sol licitation se situe a la temperature de 70Q C pour Ie stratifie ISO. A ce taux de chargement, des analyses au Microscope Electronique a Balayage (M.E.B.) met tent en evidence un endommagement du stratifie plus important dans Ie sens transversal du renfort (photo nQ 1 et 2) par rapport au sens longitudinal (photo nQ 3) au debut du fluage.

Dans Ie sens longitudinal du renfort, la rupture du stratifie s'initie par une rupture du premier pli de fibre Qil sur la face en compression de l'echantillon (PHOTO nQ 4), entrainant une chute de la contrainte a la rupture'q'" RF. Le mecanisme se propage successivement dans les differents plis du materiau, en l'affaiblissant de plus en plus jusqu'a sa rupture.

Dans Ie sens transversal du renfort, la valeur de la contrainte a la rupture depend de la mat rice et de l'adhesion resine!fibre. Dans cette configuration ou la resine est preponderante la deformation de I' eprouvette entraine une rupture des fibres (sens trame) qui ne provoque pas une chute de ~RF aussi importante que dans Ie sens longi tudinal. Ainsi, la deformation de l' eprouvette peut etre plus importante dans Ie sens transversal, ce qui explique la difference de duree de vie entre les deux sens de sollicitation.

Avec la temperature, la dilatation plus elevee de la matrice polymerique par rapport a la fibre de verre, genere des contraintes interfaciales qui modifient Ie transfert de charge resine!fibre et les caracteristiques mecaniques du composite. Ce phenomene peut expliquer la modification du facies de rupture, avec la temperature, d' une eprouvette sollicitee en flexion 3 pannes.

556

En immersion a la temperature ambiante, la rupture du s ' amorce sur la face en traction r.) ~8 S ' ini tie sur la compression lors d'un fluage en immersion a (40, 60, 80Q C)

Pour de hauts taux de contrainte (70 a 90 %~RF)

materiau face en

I' endommagement initial du stratifie est predominant vis a vis des differents processus de degradations physico-chimiques du materiau (plastification ou hydrolyse de la resine). 11 apparait donc tres aleatoire de prevoir Ie comportement a long terme d'un materiau composite pour de bas taux de contrainte (30 %"if"RF) a partir de resultats obtenus pour des contraintes elevees, en utilisant la relation suivante :~/tRF = A - B Log tR c::r/~R et tr representent respectivement Ie taux de contrainte applique et Ie temps a la rupture. La non linearite de la courbe~f(R = f(tR) est mise en evidence sur la figure 5.

Des estimations de la duree de vie ont ete effectuees a partir d'un modele Arrhenien pour les deux stratifies ISO et TETRA, en immersion a 20Q C (tableau 2).

sens de la sollicitation

sens II sens 1

ISO 21 ans 79 jours TETRA 4 ans 9 jours

Tableau 2 : Prevision de la duree de vie tR des stratifies ISO et TETRA a 50 % "" Rf en immersion a 20 Q C.

557

IV CONCLUSIONS

L' application d' une contrainte mecanique sur un materiau composite provoque un endommagement initial au sein du materiau et une elevation de la quantite d'eau absorbee a. la saturation.

Des tests de fluage en flexion 3 pannes, realises en immersion en eau distillee a. (40aC, 60ac, Boac) mettent en evidence Ie meilleur comportement du stratifie polyester isophtalique vis-a.-vis du stratifie a. base de resine polyester tetrahydrophtalique - Sa duree de vie est. plus elevee pour Ie stratifie ISO, alors que sa vitesse de fluage E est plus faible vis-a.-vis du stratifie TETRA. Cette difference de comportement est vraisemblablement due au comportement plus hydrophile et a. la plastification plus rapide de la resine TETRA par rapport a. la resine ISO.

Les experiences de fluage, effectuees a. de bas taux de contraintes (30 %fRF), en immersion a. hautes temperatures, font predominer l'effet des processus de degradation physico-chimique vis-a.-vis de l'endommagement du materiau (fissuration, decohesion ... ). Cependant une elevation de temperature peut generer d' autres mecanismes comme des contraintes interfaciales qui perturbent Ie transfert de charge du materiau.

L' etablissement des lois de comportement a. long-terme passe donc a partir d' essais de vieillissement accelere a etablir un compromis contrainte, temperature et duree de l'essai.

558

BIBLIOGRAPHIE

19th International ship structures congress, report of commitee IIIe - Genova 1985.

2 B. DEWIMILLE : "vieillissement hygrothermique d'un materiau composite fibre de verre, resine epoxyde" E.N.S.M.P, rapport de these - Juillet 1981.

3 J.N DEWAS, J.P FAVRE - Influence de cyclage (Tc, HR %) sur la cinetique de diffusion d'eau dans les materiaux composites. JNC5, Paris, Septembre 1986 380-393.

4 Ch. MELENNEC, R.JACQUEMET, A.LAGRANGE. Influence d'une contrainte mecanique permanente sur la cinetique de diffusion de.l'eau dans un stratifie polyester verre E.-rapport interne IFREMER 88/DIT/EQUEM/R09 - Avril 1988.

5 R.JACQUEMET, Comportement sous charge mecanique en environnement marin de materiaux composites a usage naval. IFREMER, rapport de these (a paraltre).

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THERMAL FATIGUE OF CARBON FIBRE I BISMALEIMIDE MATRIX COMPOSITES

T. JENNINGS, D. ELMES·, D. HULL

University of Cambridge - Department of Materials Science and Metallurgy Pembroke St. - CB2 3QZ CAMBRIDGE - England

·BP Research Centre Sunbury on Thames - MIDDLESEX - England

ABSTRACT

The design of a new thermal fatigue unit based on semiconductor thermoelectric devices to provide rapid and programmable thermal cycling is described, The unit has been calibrated to measure the temperature-time response of a typical laminate to establish the optimum conditions for short cycle times, The effects of cycling between -50°C and +SO°C on a range of cross-ply carbon fibre/BHI laminates has been studied. A qualitative analysis of the effects of thermal cycling on three [0 3 /906 1. laminates is presented.

INTRODUCTION

The availability of polymer composites capable of withstanding service conditions up to 300°C for hundreds of hours is increasing. Bismaleimide (BHI) and polyimide resins show a greater temperature tolerance than epoxy resins and so find applications in aircraft components close to the engines such as cowls, ducts and nacelles as well as in space applications. A major service requirement is the ability to withstand thermal cycling from the maximum operating temperature to as low as -55°C for aircraft and -150°C for space.

The mismatch in thermal expansion coefficient (TEC) between the longitudinal and transverse plies of a cross-ply laminate gives rise to tensile thermal stresses perpendicular to the fibre direction and compressive thermal stresses parallel to the fibre direction. The TEC mismatch arises from the different TECs of the fibre and matrix. Vith the increasing cure temperatures involved in the fabrication of bismaleimide and other polyimide systems, these stresses become significant at sub-zero temperatures, leading to the formation of transverse matrix cracks parallel to the fibres.

564

Cracking leads to a lowering of the residual properties. Thus, for example Owens and Schofield /1/ have shown a 34% reduction in the tensile strength of a [~45°)6 PMR-15 laminate after 5000 cycles from -18°C to +232°C.

Thermal cycling with ramp rate control is time consuming, /2, 3 & 4/ and typical cycles using conventional apparatus may take up to 100 minutes. This paper describes a new thermal cycling unit for rapid and controllable testing and gives experimental results for a carbon fibre/BMI resin composite system so as to establish some of the main characteristics of thermal cracking.

I - MATERIAL AND EQUIPMENT FOR THERMAL CYCLING

In selecting materials for this programme, PMR-15 and various BMI's produced in prepreg form by BP Advanced Materials were compared for microcracking as described by Elmes and Gilbert /5/. A commercial BMI system manufactured by BP Advanced Materials' US Polymeric was selected as representitive of a high Tg material susceptible to thermal micro­cracking during prolonged or severe cycling. Laminates were autoclaved by US Polymeric from unidirectional tape following recommended cure and postcure cycles. A series of [On/90m). symmetric laminates (30 cm by 30 cm) were prepared.

To give rapid but controllable thermal cycling, a powerful heat source and sink that can be kept in intimate contact with the specimen is needed. The thermal cycling units in this work /6/ use semiconductor thermoelectric "Peltier" devices which work on the same principle, but in reverse, as a thermocouple. A DC source is applied to the couple causing heat to be rejected at one junction and absorbed at the other. If the current is reversed the flow of heat is reversed and so a supply of DC current of alternating polarity results in each junction cycling above and below the ambient temperature.

A schematic diagram of the thermal cycling unit is shown in Figure 1. The laminate is placed between two thermoelectric devices, allowing heat flow from each surface of the laminate to its centre. The back face of each device is in contact with a water cooled heat sink. Yhen the laminate is cooled, heat is rejected from this face. The DC supply is controlled by a Eurotherm 818 multistage programmer allowing thermal cycles with specified ramps and dwells to be followed. The size of the specimen and the temperature range depend on the nature of the thermoelectric device.

11- CALIBRATION OF THERMAL CYCLING UNITS

During a thermal cycle the surface temperature of the thermoelectric device is measured using a copper-constantan thermocouple located at the centre of the face and held in place by a 5 cm by 5 cm square copper clamping plate. The thermocouple is connected to the Eurotherm 818 controller which displays the device surface temperature. During the heating and cooling stages of the thermal cycle a temperature gradient exists through the thickness of the laminate. The through­thickness temperature profile will depend on experimental conditions such as the rate of change of temperature and on material properties and geometry.

565

The initial mode of damage during thermal cycling of cross-ply laminates is normally matrix cracking of the inner transverse plies. This occurs at or near the minimum temperature of the thermal cycle, where the tensile thermal stress reaches a maximum. The calibration work aimed to determine the dwell time to eliminate any temperature gradient across the laminate at the peak cycle temperatures. This was studied using thermocouples placed at the surface and in the centre of a laminate. Furthermore, as the programmer was controlled from a thermocouple placed on the laminate surface, the actual peak temperatures seen by the centre of the laminate when isothermal were determined for a seven stage programmed cycle between -50°C to +50o C with a maximum ramp rate of 50°C per minute, as shown in Figure 2.

To ensure a good thermal and physical contact between the test laminate and the copper plate, a ZnO filled silicone heat sink compound was smeared onto both surfaces of the laminate. A typical experimental configuration is shown in Figure 3.

The graphs in Figures 4(i) and 4(ii) show the temperature-time response at the surface and mid-thickness of the laminate for the upper and lower peak temperatures respectively. The maximum and minimum temperatures of the laminate during a typical thermal cycle were +54°C and -48°C respectively with corresponding dwell times of 210 seconds. This gives a total cycle time of 692 seconds whilst retaining control of the ramp and dwell profiles.

The discrepancy between the actual peak temperatures in the centre of the laminate and the programmed peak temperatures of +50°C and -50°C can be attributed to the control thermocouple being at the sample surface and some temperature overshoot associated with characteristics of the Eurotherm 818 controller. This leads to a temperature oscillation. The actual peak temperatures of the laminate are assumed to correspond to the temperature at which through-thickness thermal equilibrium first occurs.

The copper plate between the thermoelectric device and the laminate acts as a thermal damper to the temperature oscillation, so the laminate temperature increases at a steady rate. The temperature at the mid-thickness of the laminate lags behind that at the surface. If the heat plates were held for a longer period of time then thermal equilibrium would eventually occur across laminate at +50°C and -50°C respectively. The main experimental advantage of this new thermal fatigue unit is to enable rapid thermal cycling of laminates. For this reason, the temperature extremes seen by the laminate have been assumed to be +54°C and -48°C corresponding to the minimum times required to attain through-thickness thermal equilibrium.

III - THERMAL CYCLING OF [03 /906 1. BMI LAMINATES

Three specimens were prepared in an identical manner to that of the calibration specimen but without a hole for a mid-thickness thermocouple. The long edges were polished to a 1 micron finish for microcrack examination. Each laminate was cycled incrementally to a total of 2,500 thermal cycles and examined periodically. Laminates were removed after a whole number of cycles (position A in Figure 2) so as to prevent discontinuity in the thermal loading. A zinc iodide die penetrant was used in X-ray analysis to monitor crack development.

566

Figure 5 shows the development of damage in the three laminates. There are some significant differences between the specimens but the general trends are the same. Transverse matrix cracking occurred initially in the central (90°) plies and later in the surface (0°) plies. All the transverse cracks spanned the width of the laminate. Observation of the polished edges showed that the majority of these cracks branched close to the 0°/90° interface.

Figure 6 shows the different modes of cracking observed. Fibre-matrix debonding is a significant feature of transverse crack initiation and was observed in all three laminates in the region of the 90° plies nearest to the 0°/90° interface. The debonds had two main growth directions in the early stages of fatigue life; parallel and perpendicular to the direction of the 0° plies. Yith increasing number of thermal cycles the "perpendicular" direction debonds developed into transverse cracks, whereas the "parallel" direction debonds grew away from the laminate edge towards the centre to form an 'intralaminar crack. At later stages of fatigue life a third growth direction was observed resulting in the formation of short cracks oriented at an acute angle to the 0° plies; these cracks grew on either side of the transverse cracks. The rate of crack growth falls off with increasing number of cycles although the characteristic "saturation" crack spacing was not observed after 2,500 cycles.

The results show that transverse cracks initiate near the 0°/90° interface and grow across the thickness of the transverse plies. Figure 7 shows a thermal stress profile across the transverse ply thickness for an unrestrained cross-ply laminate at a temperature T below its stress-free temperature To' At a position far from the ends the stress profile may be assumed to be uniform across the plies as predicted by laminate analysis. The laminate is stress-free at the ends and this influences the stress profile which becomes less uniform towards the free surface (Figure 7(i». The magnitude of the tensile thermal stress in the 90° plies near a free surface therefore decreases from the predicted value, moving away from the interface towards the centre of the laminate. Overall the transverse tensile stress will increase with decreasing temperature until it exceeds the transverse ply tensile stress, initiating a mode I matrix crack. As a result of the stress profile the crack is most likely initiated near the 0°/90° interface. This is the case near a free surface, whereas away from the end a crack may initiate at any position across the transverse ply thickness. The crack itself then acts as a free surface causing a stress redistribution in the transverse plies. Subsequent cracks inititating away from the free end of the laminate are therefore more likely to initiate near the 0°/90° interface under the influence of the new free surface caused by the first transverse crack.

The initiation of fibre-matrix debonding parallel to the 0° ply direction may be explained either in terms of mode II cracking due to thermal shear, or mode I cracking due to out-of-plane edge stresses at the free surface (az and ~xz)' During parts of the thermal cycle steep temperature gradients may exist across the thickness of the laminate. This causes differential through-thickness thermal expansion of the plies. As a result thermal shear stresses may develop in the 90° plies parallel to the 0° ply direction, leading to mode II cracking. An alternative explanation lies in the consideration of the stress state at the free edge of the laminate. Pipes and Pagano /7/ have applied

567

finite difference techniques to analyze the stress state of a four­layer symmetric angle ply laminate under uniform axial extension. Their results have shown a large increase in the out-of-plane stresses az and Lxz towards the laminate edge. Such high stresses may cause mode I intralaminar cracking.

IV - CONCLUSIONS

A new thermal fatigue unit using Peltier devices has been designed which gives relatively short cycle times. It has been used to investigate the development of damage in carbon fibre/BMI laminates. Characteristic crack growth morphologies have been identified and their relationships with the stresses arising during thermal cycling is discussed.

v - ACKNOYLEDGEMENTS

The authors wish to thank the British Petroleum Company for permission to publish and BP Advanced Materials (US Polymeric) for the supply of materials.

REFERENCES

1. Owens G A, Schofield S; Comp Sci & Tech, 33(3) (1988) 177-190. 2. Adams D S, Bowles D E, Herakovich C T; J Reinf Plas & Comps, 5 (1986)

152-169. 3. Tompkins S S, Yilliams S L; J Spacecraft, 21.3 (1983) 274-280. 4. Fahmy A A, Cunningham T G; NASA-CR-2641 (1976). 5. Elmes D A, Gilbert D G; Proc Fibre Reinf Composites '88, Liverpool

(1988) 36/1- 36/10, Publ Plast & Rubber Inst. 6. UK Patent Application, GB 2188 163A, 23/9/87. 7. Pipes R B, Pagano N J; J of Comp MatIs, 4 (1970) 538-548.

568

,0 )8

56

;2

50

'8

'6

'4 '2 '0 0

3 1 2

4

5 1 2 3 4 5

Specimen Copper plate Thermoelectric Heat sink Coolant Flow

Fig.l - Schematic of thermal cycling unit

Dwell o 50·~-----------------.r---~---. II ...

30 · ~~ ______________ ~ ____ ~~

:> 1 O·

f -1 00 II

~ - 3 0 • ..... ----~ ...... ~t".-----------~ -50 .+-_________ ....L ____________ _

Time

Fig.2 - Typical thermal cycle

Direction of heat flow

O·

90· 9i==~~~~~~~2cm o • ~5cm Thermocouple

D' t' f ;... placed in O. 4mm hole 1~:~t1~~o~ •• , at midthickness

of laminate

Fig.3 - Test specimen configuration

-30 -32

~ -34 -36 -& Av Temp. edge

! -38 ..... Av Temp. centre :>

E -40 OIl -42 ~ -44-~ -46 f- -48

-0- Av. Temp edge 8, -50 ..... Av. Temp.centre f -52

~ -54 00( -56

-58 -60

00 0000000 0 000 0000000

"'''' O)"'lnCD- ..... f'.. 0 "'''' OloC\lll)CX) ....... "If,.... _ ..... .,.... "'C\lN '" ....-....--....--NNN

Time/sees Time/sees

device

0 0 C'l

(i) - Upper peak temperature Fig.4 (ii) - Lower peak temperature

14

12 • • II ••• .>< 10 •••• • 190' u l!! • •• •••• • • 100

U 8 • ••• '0 • 290' •• i 6 •• •••• • c 200

E • • •••• • • 390' :) 4 .... cc c • 300

Z aaccc cc 2 a ..

• 0 0 0 0 0 0 0 0 0 0 0 0 <0 0 <0 0 '" 0

C\I C\I M

Number of Cycles

Fig. 5 - Damage development in three laminates

Matrix crack parallel ,,0 _t<ix cf.c. Angled crack

Fig.6 - Modes of damage

x

+

<I) O' 90 ' O'

Fig.7 - Thermal stress profile in transverse plies at T<To.

(i) Near free surface (ii) Far from free surface.

569

THE INFLUENCE OF TEMPERATURE AND MOISTURE ON CROSS·PLY CRACKING IN CFRP IN TERMS OF MATRIX

FRACTURE STRAiN AND INTERFACE STRENGTH

ABSTRACT

P.w. M. PETERS, S.1. ANDERSEN'

DFVLR - Institut fOr Werkstoff-Forschung PO 906058 - 5000 KOLN 90 - West Germany

*Riso National Laboratory Postbox 49 - 4000 ROSKILDE - Denmark

Cross-ply cracking in specimens out of the laminate 0/90 4/0 was investigated in dry condition and after satu­ration with water at RT (moisture uptake 1.44 %) in the temperature range of -100°C to +100°C.

These investigations resulted in Weibull transverse fracture strain distributions, with which the strain at first ply failure (FPF) and at interface failure was de­termined. At increasing temperature in the dry material two conflicting phenomena influence transverse cracking: these are the ductility of the matrix and the strength of the fibre-matrix interface. The improvement of ductility increases the transverse fracture strain, whereas the re­duction of the interface strength reduces the transverse fracture strain. For the strain at FPF the improvement of ductility dominates, except in the range between RT and 60°C where the interface strength drops considerable. Saturation with moisture at RT and testing at RT leads to a further improvement of ductility and an additional de­crease of interface strength. The overall effect on the strain at first ply failure is positive. The wet material tested at -100°C showed both an increase in strain at FPF and at interface failure.

INTRODUCTION

With the increasing use of graphite/epoxy materials for structural applications it is necessary to study the

572

mechanical behaviour of these materials under different environmental conditions. Matrix and interface dominated properties are expected to vary significantly as a function of temperature and moisture. The transverse strength or transverse fracture strain in an angle ply laminate for example is such a mechanical property. If the influence of temperature and moisture on laminae instead of laminate properties is considered, the residual strains play an important role. Thus the aim of the present investigations is twofold:

1. to determine the residual strains as a function of the test temperature and moisture content

2. to determine the dependence of the transverse fracture strain as a function of the temperature and moisture content.

The transverse fracture strain is investigated with the aid of the laminate 0/90 4/0 out of the system T800/ R6376 of the company Ciba-Gelgy. This system contains the intermediate modulus T800 carbon fibre of Torayca company. The way to describe the transverse fracture strain or strength can roughly be divided into two methods:

the fracture mechanics approach /1-4/ the statistical approach /5-8/.

In the present investigation the statistical approach will be applied. This approach is based on the assumption that fracture in the 90°-ply occurs as a result of distri­buted defects of different size. The laminate with the clustered 90°-plies 90 4 is chosen, because in this lami­nate multiple fracture can be measured easily with the aid of a piezo-electric transducer in connection with a computer /6-8/. Based on these investigations a model was developed /9/ which describes the transverse fracture strain as influenced by the defects in the material, the ductility of the matrix, the constraining effect of neigh­bouring layers and the strength of the interface. This model is schematically presented in Figure 1. The fracture strain distribution in the ideal case is presen­ted by a perpendicular line (no defects, thus no scatter). In reality, materials always contain defects which cause the fracture strain to decrease. Dependent on the distri­bution of defects the fracture strain distribution shifts to lower fracture strains under a certain angle a (a = shape parameter).

There are two effects which counteract this decrease of fracture strain. These are the ductility of the matrix and the constraining effect exerted by the neighbouring layers. Both effects reduce the severity of the defects and thus displace the fracture strain distribution to

573

In the present investigation the fracture strain dis­tribution of the 90 D-ply in a 0/904/0 laminate is deter­mined in the temperature range of -100 DC to +100 °C for dry material and for material which was saturated with water at RT. The results are interpreted with the above mentioned model.

I MATERIALS AND EXPERIMENTS

The laminate 0/90 /0 was produced in a press-clave (a press with a 4 cm high press chamber between the heat­ing plates) at 175 DC and at a pressure of SOO kPa in 2.5 h. Vacuum was applied during the heat-up phase of 3 DC/min until a temperature of 140 DC was reached. An ultrasonic inspection of the laminate (C-scan) did not show any voids. An additional unidirectional laminate Os was produced to determine the elastic constants EN and EL and the thermal expansion coefficients ap and a~. In order to be able to determine the thermal strains in the 0/904/0 specimens at the different testing temperatures, the ~hermal expansion coefficients an and aL were measured as a function of the temperature with a dilato­meter.

From the produced plates 16 rom wide and 200 rom long specimens were cut with a diamond wheel. In the gripping area, 50 rom long aluminium tabs were glued on the speci­men, which results in a specimen gage length of 100 rom. These specimens were loaded up to fracture in a displace­ment controlled (v = 2 rom/min) Instron testing machine. Tests on dry specimens were performed at -100 °C, -40 °C, RT, +60 DC and +100 °C respectively. The experiments below RT were performed in a container with alcohol which was cooled on the outside with the aid of liquid nitrogen. The experiments in the positive temperature range were per­formed in water heated with the aid of a heating element. To prevent an influence of the medium on the fracture be­haviour, the specimens were wrapped in aluminium foil.

The wet specimens are from plates produced for ear­lier investigations /S,9/. A slightly different cure cycle was applied for this material (pressure 600 kPa). The spe­cimens were saturated in destilled water for period of about 11 months. The specimens were found to be saturated after ~ 3 months at a moisture uptake of 1.44 % by weight. The specimens were stored under water for a longer period of time because of experimental problems and the avail­ability of the testing machine. The submersed specimens had no tabs. Application of end tabs with the aid of a room temperature adhesive after saturation proved to be impossible, because of the low strength of the bond line. For this reason these specimens were tested with no tabs, which caused in some cases premature specimen failure at

574

a low stress. Because of the limited number of specimens available, some failures during preliminary testing and completely suppression of transverse cracking at +60°C and 100°C for the saturated material only results at -100°C and at RT can be presented. The tests at -100°C were performed as for the dry material with the specimens wrapped in aluminium foil (after they were saturated with water) .

Multiple fracture of the 90°-ply was measured with the aid of a piezo-electric transducer in connection with a computer. The dynamic load change as a result of crack­ing of the 90°-ply is thus registered with the actual value of the load and the strain. The strain is measured with the aid of a strain gage in the middle of the speci­men.

II RESIDUAL STRESSES

The residual thermal strain can be calculated with the aid of the laminate theory based on the thermal expan­sion coefficients of the different plies. For the present investigation the thermal expansion coefficients parallel and transverse to the fibre direction have to be deter­mined, which can be done either by a strain gage technique /10/ or by a dilatometer. A complete deduction of the ther­mal residual strain in the transverse plies based on the measured thermal expansion coefficients and considering the coupling expansion is presented elsewhere /11/. The results for the different temperatures are presented in Table 1 and Figure 2. The swelling behaviour due to water uptake is investigated with the aid of an unbalanced O2/90 7 laminate, which shows after cooling down from the production temperature a curvature due to the mismatch of the thermal expansion coefficients in and perpendicular to the fibre direction. The curvature of 3 rom wide spe­cimens at RT before and after saturation with water (by submersion in destilled water) was measured. From the curvature of these specimens the thermal expansion coeffi­cients as well as the swelling coefficients can be deter­mined making use of the theory of Timoshenko /12/ derived for a bimetallic strip.

The swelling strain in the transverse ply of a 0/90 /0 laminate can (similar to the thermal strain /6/) be deteF­mined with the simple equation:

ErI'I= -C (13 2 - 13 1 ) E1 (1) 2.

n/2 • E2 + E1

where 13 1 , 13 2 are the swelling coefficients in and trans­verse to the fibre direction and C is the moisture con­tent. From the Timoshenko equation C(132-131 ) was determined

575 -3 . to be 5,857 x 10 m/m at complete saturat~on (Cm

1.44 %). Thus the difference in swelling coeffic~8~t meas­ures 32 - 31 = 0.406. Substitution of the values of the respec~ive constants in eq (1) leads to E~ = 0.427 %. Now the total residual strain is considered to be

(2 )

assuming that the swelling does not depend on the tempera-ture (E%m -0.427 % at all test temperatures). The re-sulting residual strain in the transverse ply is indicated in Table 1 and Figure 2. Heating or cooling of the medium in which the test specimens were submersed up to the re­quired test temperature was done as fast as possible and the temperature of the medium (alcohol or water) was kept constant for 15 min. Thus the moisture content can assumed to remain unchanged, especially in the inner transverse plies.

III DETERMINATION OF THE TRANSVERSE FRACTURE DISTRIBUTION FROM THE DATA OF A SINGLE SPECIMEN

The analysis which enables us to determine the trans­verse fracture strain distribution based on the data of multiple cracking in the 90 0 -ply is extensively treated elsewhere /6,7,8,9/. For this reason the analysis is only presented briefly.

The occurrence of multiple cracking makes it reason­able to consider the 90 0 -ply to consist of a number of elements all of which can break. The number of elements in the 90 0 -ply is determined with the aid of the stress distribution in the 90 0 -ply close to an existing crack. This stress distribution can be calculated with a shear lag analysis /13,14/. The stress in the 90 0 -ply (ply 2) as a function of the distance x to the crack determined this way is given by:

°2, x = °2, 00 [1 - exp ( '1x) 1

where ° 00 is the undisturbed stress at infinity x = 00

(thermat plus mechanical stress) and Y is given by

'( = JG/b (11 (E1 a 1) + 2/(E2 a 2 ))

(3 )

(4 )

in which b is the width of the shear transfer layer and a] and a 2 are the thickness of the surface and 90 o -plies respectively.

The shear lag parameter G/b is the main difficulty in applying the shear lag model /13/. In the present investi­gation the same procedure as in /6-8/ was followed. G is assumed to be the shear modulus of the matrix whereas b

576

is taken to be twice the fibre diameter (b = 0.010 rnrn). At a short distance to an existing crack the stress in the 90°-ply according to Eq. (3) is small so that in this range of low stress no more cracking will occur. This leads to the often found regular crack spacing in the 90°-ply. Thus the element length should be chosen suffi­ciently large to enable the possibility of fracture in every element. The element length, however, should not be chosen too large to prevent the possibility of multiple cracking inside a single element. This would occur if the element length is chosen to be the length in which the stress in the 90°-ply almost completely recovers. For this reason the element length I is somewhat arbitrarily (similar to the early work of Ro~en /15/ for the ineffec­tive length of a broken fibre) chosen such that at its borders 90 % of the undisturbed stress is reached. Now the model implies that the stress has a simplified step function. The stress in a broken element is zero, whereas in the neighbouring elements the stress is undisturbed (07. .. 00)' With the aid of the known element length I the number of elements N in the specimen is determined~

For a statistical description of the fracture strain use is made of a two-parameter Weibull distribution given by

F = 1 _ exp _ (:L IQI a (5)

Ef,ld which expresses the probability of failure F of an element (length I ) as a function of the strain Ef 1 (= E2 00)' The two W2ibull parameters are the shape par~meter a and the characteristic fracture strain fr 1 (substitution of Ef = E' 1 leads to F = 0.632). TniSoWeibull distribu­t~6~ocan b~ ~resented as a straight line if ln (-In(l-F)) is presented as a function of ln E . The two Weibull parameters define the position of th~Ocurve: a is the slope of the line, whereas the characteristic fracture strain is the strain at F = 0.632. Now from the available experimental data (list of occurred cracks with respec­tive fracture strains in increasing order of strength) the Weibull distribution is determined. For every data point the probability of failure is approximated by

j/(N + 1) (6 )

where j is the crack order number, whereas N is the number of elements in the specimen. The data are now transformed in couples of ln E and ln (-In(l-F~-values and with the aid of linear regr~s~ion and least square analysis a best fit Weibull distribution is determined. If a Weibull distri­bution for elements with length 1 is thus determined, the o

577

fracture strain distribution of any other ply-thickness or element length (or specimen length) can be calculated with the aid of

(7 )

which describes the influence of the volume (V , V ) on the characteristic failure strain (the shape pAram~ter is unchanged) according to the Weibull theory. This equation is used to determine the characteristic fracture strain at first ply failure in 100 rnrn long specimens.

Figure 3 schematically presents a Weibull distribu­tion, drawn uninterrupted in the range of the experimen­tal data and drawn intermittent for the extrapolation to higher probabilities of failure. It is clear that first ply failure is caused by the biggest flaw (as indicated schematically in Fig. 3). Moving on the element fracture strain distribution curve to its top, fracture of the re­spective elements is caused by smaller and smaller defects. If we assume that the last breaking element is free of defects (which can only be said of macroscopically and microscopically void free material), then this fracture strain can be accepted as an interface failure strain, or in case the interface failure strain is stronger than the matrix, as an ultimate matrix strain. The gage length of the specimen being 1 = 100 rnrn contains roughly 100 to 200 elements and thus the strain at interface failure for the strongest element can be calculated by substituting the corresponding value of F = 0.99 to 0.995 (= j/N+1 = 100/101 or 200/201) in eq. (6).

IV WEIBULL FRACTURE STRAIN DISTRIBUTIONS AT DIFFERENT TEST TEMPERATURES FOR THE DRY AND WET MATERIAL

Multiple cracking was measured in 10 dry specimens at the different test temperatures and in 5 specimens sa­turated with water at RT and -100°C respectively. Before fracture strain distributions can be determined first the element lengths at the different environmental conditions are calculated.

The element length 1 of the 90°-plies at the respec­tive testing conditions i~ calculated with Eq.(3) after substitution of 02 (x = ± 10/2) = 0.9 02 00' The element length 1 depends b~th on the temperature as well as on moistureOthrough E2 and the shear modulus G of the matrix. In dry condition tne dependence of E2 on the temperature follows from Figure 4 whereas the dependence of G (also indicated in Figure 4) on the temperature was taken from the literature /16/. From the variation of the resin ten­sile modulus with temperature (3.6 GPa at room tempera­ture and 2.5 GPa at 120°C /16/) a linear dependence of the shear modulus G was calculated.

578

The influence of humidity on the elastic constants E and G can be determined only roughly because of the s~all number of tests performed. The transverse stiffness can expect to decrease by 7 % at RT if the material is saturated with moisture as measured for different other CFRP materials /17/. The dependency of the variation of the transverse stiffness of the wet material on the tem­perature was determined by comparing the E-modulus of the wet specimens out of the laminate measured at different temperatures. The change in E-modulus of some particular specimens determined at different temperatures was com­pletely contributed to the change in transverse stiffness. Thus the following relation between the transverse stiff­ness E;and the temperature

(8 )

It1 * w~th the transverse ~odulus at OOC E,' = 8.634 GPa and C = -0,01465 m/m/oC was found. The shearing modulus was assumed to decrease at room temperature with the same amount ( 30 %) as the E-modulus of the pure resin 914 C under the influence of moisture uptake upto the satura­tion level/10/. The variation of the shearing stiffness with the temperature was further assumed to be identical to the case for dry material. The resulting dependencies of the moduli for the saturated material are indicated in Figure 4.

The resulting element lengths at the different test­ing conditions are indicated in Table 1. The available crack data are treated as described in the analysis and presented in Table 1. A complete description of the test results for the dry materials is published elsewhere /11/. In most cases the crack data can be described by a two­parameter Weibull distribution as presented by the straight lines. For higher temperatures, however, at a larger prob­ability of failure the data deviated from the initial li­near relation. For this reason a best two-parameter Weibull fit is determined in the initial about linear range. In Table 1 is indicated the best fit weibull distribution and the number of cracks which are considered for the Weibull fit. Figure 5 presents the crack data and the determined weibull distributions for the RT-test on the dry and the saturated material.

V DISCUSSION

The determined Weibull distributions can be interpre­ted with the aid of the model explained in Figure 1 of the introduction. It shows how the transverse fracture strain is influenced by the defects, the ductility of the matrix, the constraining effect and the strength of the interface. Let us consider now, what parameters vary in the present

579

investigation. Under the influence of moisture and varying testing temperatures the mechanical properties of the resin and the interface change.

5.1 THE INFLUENCE OF THE TEMPERATURE

The fibre-matrix bond is established at relative low temperatures compared with the curing temperature /19/. Thus it can be expected that at higher temperatures the fibre/matrix interface degrades before the matrix does. The fracture strain of the matrix, however, increases with increasing temperature. Thus there are two conflict­ing influences on the transverse fracture strain. At in­creasing temperature the increase of matrix fracture strain tends to increase, whereas the reduction of the interface strength tends to decrease the transverse frac­ture strain. The influence of these conflicting effects on first ply failure and on the strain at interface fail­ure is made clear in Figure 6. It shows that at increas­ing temperature the strain at first ply failure increases exept in the range between RT and 60°C. This indicates, that the positive influence of the increasing matrix duc­tility dominates, except in the range between RT and 60°C, where the strength of the interface drops considerable.

5.2 THE INFLUENCE OF MOISTURE AND TEMPERATURE

Moisture generally has the same effect as the tempe­rature, it reduces the strength of the interface and im­proves the ductility of the matrix. This is confirmed by comparing the test results of the tests at RT of the dry material with those of the material saturated with mois­ture (Figure 5). Two aspects show that the wet material has a more ductile matrix, this is the shape parameter (a = 16.14 for the wet material compared with a = 11.14 for the dry material). Further the strain at FPF (see Figure 6) for the wet material is larger than for the dry material. The strain at interface failure, however, is lower than for the dry material which could be expected as mentioned before.

The test results at -100°C for the material saturated with moisture compared with the results for the dry ma­terial show an increase for the strain at tirst ply fail­ure as well as for the strain at interface failure. The increase in strain at FPF can be explained by the in­creased ductility of the matrix. The other aspect of in­creased ductility, an increase in the shape parameter, is not found. This could be caused by the somewhat unexpec­ted increase in interface strength for the wet material. As these findings are based on only a few tests further investigations are necessary to study the influence of moisture at low temperatures. Further, it has to be clari-

580

fied whether the assumed independency between thermal ex­pansion and swelling is valid.

CONCLUSIONS

The thermal strain in the transverse plies of the 0/90 /0 laminate was calculated based on the measured ther~al expansion coefficients of the unidirectional~ laminate (an and a~). This thermal strai~ measures E? = 0.218 % at +100 °c and increases up to E = 0.665 % ~t -100°C. The swelling strain due to satufation with moisture (moisture content 1.44 %) was determined E: -0.429 %, so that the saturated specimens at -100°C and +100oC have a residual strain in the 90 -plies of -0.211 % and 0.236 % respectively. Tensile !ests were performed under the different environmental conditions and multiple cracking was studied. As a re~ults of the negative residual stresses in the saturated specimens transverse cracking did not occur at 60°C and +lOO°C. It was shown, that the fracture strain (mechanical + resid­ual strain) distribution determined from the crack data of a single specimen could be described by two parameter Weibull distributions.

A model developed previously /9/ proved to be able to interpret the different weibull distributions. Two phenomena play a dominant role in the dependence of the transverse fracture strain on temperature and moistur~. These are the fracture strain and ductility of the matrix and the strength of the interface. At increasing test tem­perature and due to moisture the matrix fracture strain improves, whereas the interface strength is reduced. From the determined Weibull distributions for the dry material it could be concluded that the strain at first failure in the 90°-ply (FPF) is dominated by the positive influence of the increased matrix fracture strain at increasing tem­perature. An exception occurs in the range RT to 60°C, where the strong reduction of interface strength over­rules the positive effect of the increased matrix frac­ture strain. After saturation with moisture the Weibull distributions determined from specimens tested at room temperature indicate moisture improves the strain at FPF (through the increased ductility of the matrix) but re­duces the strain at failure of the interface. The some­what unexpected results at -100°C show both an increase in strain at FPF as well as for the strain at interface failure. This has to be verified and explained by addi­tional experiments.

REFERENCES

/ 1/ Parvizi, A., Garret, R.W. and Bailey, J.E., J. Mater. Sci. 13 (1978) pp 195-201

581

/ 2/ Flaggs, D.L. and Rural, M.H., J. Composite Mater. 16 (March 1982) pp 103-116

/ 3/ Crossman, F.W. and Wang, A.S.D., ASTM STP 775 (American Society for Testing and Materials, 1982) pp 118-119

/ 4/ Nuismer, R.J. and Tan, S.C., in: Mechanics on Compo­site Materials, Proc. IUTAM Symp. on Mechanics of Composite Materials VPI, Blacksburg, VA, USA, 1982

/ 5/ Manders, P.W., Chou, T.W., Jones, F.R. and Rock, J.W. J. Mater. Sci. 18 (1983) pp 2876-2889

/ 6/ Peters, P.W.M., J. Compo Mat., Vol. 18 (1984), pp 545-557

/ 7/ Peters, P.W.M., ASTM STP 907 (1986), pp 84-99

/ 8/ Peters, P.W.M., Meusemannn, H., ICCM-VI, ECCM II, London (1987), pp 3.508-3.525

/ 9/ Peters, P.W.M., ASTM STP 1012 (1988)

/10/ Carlsson, L.A., Pipes, R.B., in: Experimental Characterization of Advanced Composite Materials. Prentice Hall, New Yersey, 1987.

/11/ Peters, PWM., Anderson, S.I., to be published in J. Compo Mat.

/12/ Timoshenko, S., J. Optical Society of America, 11, (1925), pp 233-255

/13/ Highsmith, A.L., Stinchcomb, W.W. and Reifsnider, K.L., VPI-E-81.33. Dept. of Engng. Sci. and Mechanics, virginia Polytechnik Inst. and State University, Blacksburg, VA, USA, (Nov. 1981)

/14/ Peters, P.W.M., Poster paper presented at Conf. on Test­ing Evaluation and Quality Control of Composites, Univer­sity of Surrey, Guildford, UR 13-14 September 1983

/15/ Rosen, B.W., AlAA Journal, Vol. 2, No. 11, Nov. 1964, pp 1985-1991

/16/ Ciba Geigy, 6376 - Data sheet FTA 140e, (1984)

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ARTIFICIAL AGEING OF FIBRE REINFORCED COMPOSITE MATERIALS· THREE STAGE METHOD

N. MARKS, A. DALZIEL

Westland Helicopters Ltd BA202 YB YEOVIL SOMERSET - England

ABSTRACT

Yhen Yestland Helicopters realised the potential of epoxy based composites for structural applications, a programme of work was instigated to investigate the properties of aged composites. After some initial work on ageing specimens artificially, a three stage technique was adopted to accelerate the ageing process. This method was chosen because it would result in a specimen with a consisteRt moisture content throughout the laminate. The conditions required to achieve this were predicted using a computer programme but it was found that the empirical results differed considerably from those predicted because the matrix system employed for these laminates does not follow Ficks laws of diffusion at higher temperatures.

INTRODUCTION

During the 1970s, Yestland Helicopters realised the potential of composite materials for manufacturing rotor blades for helicopters. An initial programme of work indicated that a 125°C curing modified epoxy, Fibredux 913, was the most suitable matrix with E glass and a PAN based carbon as the reinforcements. The majority of the materials employed would be in the form of a 300 mm wide unidirectional tape.

Although it had been known for some time, that epoxy resins absorb moisture, it was not until the late 1970s that it was appreciated that this could have a significant effect on the mechanical properties. It was at this time that Yestland

588

Helicopters, in conjunction with RAE Farnborough, instigated programmes of work to investigate the effect of both artificial ageing, in humidity cabinets, and natural ageing on the properties of composite materials. This paper examines some of the work carried out on artificially aged specimens. Although the ultimate aim of the work was to determine the effects on the strength of the materials, it is intended to deal only with the moisture management section of the programme.

INITIAL AGEING

Our initial work concentrated on environmental ageing under fixed conditions: 35°C, 95 to 100% RH, until the specimens attained a known moisture content, which was calculated as being approximately 80% of the equilibrium moisture content (Moo). This work proved that ageing affected the mechanical properties but it was not known whether this was due to (a) pressence of moisture (b) stress gradient caused by the through thickness moisture gradient or (c) a combination of the two. It was then decided to examine alternative methods of obtaining aged specimens to produce specimens with a constant through thickness moisture profile.

SPECIMEN MANUFACTURE

Laminates were produced from Fibredux 913 preimpregnates using a press-clave technique for consolidating under a pressure of 550 kPa. It was necessary to produce laminates with a void content of less than 2%. This was controlled by a non destructive examination of the complete laminate using ultrasonics with c-scan, a visual examination of a cross section and determination of the fibre volume content.

MOISTURE MANAGEMENT

The conditions under which specimens were to be aged were determined, with the close cooperation of RAE. In determining these conditions, a computer program, produced by RAE and known as DIFF 41 ,2 was employed. This program has been widely used for predicting the behaviour of epoxy composites in hot wet environments. It is based on the presumption that moisture diffuses into an epoxy composite structure in accordance with Fick's Laws of Diffusion.

The method chosen to obtain specimens with a consistent through thickness moisture profile in a relatively short time required three separate conditioning stages.

All specimens were initially dried in a vacuum oven.

589

Stage 1 Accelerated ageing stage

This stage enables the moisture to diffuse rapidly into the centre of the laminate. To achieve this, high temperatures and humidities are used.

Stage 2 Drying stage

The high moisture content at the surface which results from stage 1 is removed and moisture diffuses rapidly towards the centre of the laminate by heating the specimens in a vacuum oven at 35°C.

Stage 3 Equilibrium stage

For this stage lower temperatures and/or humidities are used than those employed at Stage 1. Some of the surface moisture is replaced whilst further diffusion occurs into the centre of the laminate until an equilibrium or "quasi-steady-state" is reached for the particular environment.

This process is illustrated in Figure 1.

For our particular programme, three different stage 1 environments were employed, all at 75% RH, with temperatures of 35°C, 45°C and 60°C.

Similarly three stage 3 environments were employed, all maintained at room temperature but at RH levels of 23%, 42% and 55%.

DEFINITIONS

1. 2.

3.

4.

Temperature of the environment T - units °C Relative humidity of the environment RH - expressed as a percentage. 2 Diffusion coefficient D - units mm per second. For a given material and thickness, D depends chiefly upon T and governs the rate of moisture diffusion into the laminate. Equilibrium moisture content Mm - this is achieved in a laminate after exposure to a particular environmental condition (T, RH) when there is no longer a net influx of moisture. The laminate could be said to have reached saturation under these conditions and there then exists a state of dynamic equilibrium or "quasi-steady-state". It is normally quoted as weight percentage of moisture in the laminate.

For a given material and thickness, Mm depends chiefly upon RH and governs the volume of moisture diffusing into the laminate.

590

TEST PROGRAMME

The test programme was based on using pretravellers to calculate the parameters required to estimate the conditions which were necessary to obtain the requisite moisture profile. The moisture uptake of the pre travellers (M) was measured after different times (t). By plotting M against t, it is possible to determine Mm and D for the laminate. It soon became obvious that this would take too long, therefore in the initial stage, the DIFF 4 program was used to estimate the stage 1 targets. For the same reason, Mm for the stage 3 conditioning was also taken from DIFF 4.

There are two main routes through DIFF 4 program, depending upon which parameters were available- for input.

(i) Input T and RH - By using RAE experimental data stored in the program, values for D and Mm are calculated for the differing environments.

(ii) Input D and Mm - Experimentally derived values for D and Mm must be provided for the particular material thickness /exposure conditions being modelled. Experimental methods and standard formulae were used to determine Mm and D using the following procedure:-

(a) Dry specimens of thickness t were exposed to conditions for temperature T and relative humidity RH.

(b l Specimens were accurately weighed after 1, 4, 9 and 16 days.

(c) Moisture content was plotted against t and the initial straight line gradient(s) evaluated.

(d) Ageing continued until curve flattens sufficiently to estimate Mm.

(e) D is estimated as follows:-

D n.s2 (d )2 where d = thickness (4Mm) of specimen

Yhen the stage was reached when it was possible to calculate Mm from the pretravellers, it transpired that the empirical values did not match the estimated values. Mm appeared to be temperature dependent as well as RH dependent -see figure 2. Since Mm on the DIFF 4 program correlated with the Mm for 60°C empirical results, it therefore follows that the original target moisture contents predicted for the end of stage 1 and hence for the end of stage 2 were too high for specimens aged at 35°C and 45°C. Mm for stage 3 was also too

high. Because of this the predicted times of exposure were also too high.

591

It was therefore necessary to amend the exposure programme. In doing this, new methods of using the DIFF 4 computer program were found. For a given Hm and laminate thickness a through thickness moisture profile independent of the diffusion coefficient can be obtained - D only controlling the time to reach that profile. For a 4efined Hm and laminate thickness the moisture level is inversely proportional to D. The time taken to reach Hm is the same for any Hm value, if the same D and laminate thickness are input.

To recalculate the exposure programme, Hm values were taken from the pretraveller results at the three different temperatures (as shown in·table 1). Hm for the stage 3 environments were calculated assuming that Hm values at room temperature were identical to those at 35°C for the same RH and that Hm has a linear relationship with RH. Predictions were based upon target moisture contents rather than exposure times so eliminating the need to calculate D. By closely monitoring weights of specimens, it was possible to transfer them to their new environments immediately they reached the predicted target weight.

The main conclusion from this work is that, for this particular resin system, Hm for a particular RH appears to be dependent upon temperature. There is, however, only a small difference between Hm at 35°C and that at 45°C. There is a much greater difference for Hm at 60°C. It must therefore be assumed that the moisture diffusion of this system does not obey Fick's laws of diffusion at temperatures of 60°C and above. It does appear to obey Fick's laws at 45°C and below. This confirms the findings of Collings and Copley2 found subsequent to the start of this work. Since there is no data for moisture uptake between 45°C and 60°C, it has been the policy of Westland Helicopters to restrict ageing to a maximum temperature of 45°C.

592

MATERIAL THICKNESS EXPOSURE TEMPERATURE mm

35°C 45°C

CFRP 1 2.8 3.1 2 2.5 3.0

2.5 2.4 3.0

GFRP 1 2.3 2.8 2 2.4 2.9

2.5 2.5 3.0

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TABLE 1

Equilibrium moisture contents achieved in pretravellers conditioned at 75% RH

60°C

3.9 4.1 3.7

3.7 4.1 4.4

3.8 2.9

(all figures expressed as percentage weight in resin) .

Acknowledgement to MOD R.A.E. Farnborough for sponsoring this work.

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REFERENCES

1 - Copley S, Computer program to model moisture diffusion RAE Tech Report TR 82010

2 - Collings T, Copley S, On accelerated ageing of CFRP, Composites Vol 14 No 3 July 1983

ASPECTS OF THE THERMAL DEGRADATION OF PMR-15 BASED COMPOSITES

F. JONES. Z. XIANG

University of Sheffield School of Materials. Northumberland Road S10 2TZ SHEFFIELD - England

Abstract

The kinetics of thermal degradation and its consequent effects on the mechanism of microcracking in carbon fibre composites fabricated from PMR-15, an end-capped bismaleimide resin system, have been studied under oxidative conditions. The overall degradation involves three stages. For the unidirectional laminates, oxidative decomposition of resin appears to control the micromechanics of their durability, whereas changes in thermomechanical properties of the matrix seems to be responsible for the formation of transverse cracks in the cross-ply laminates.

1. Introduction

Thermo-oxidatively stable polymer systems which overcome the limitations of conventional epoxy resins to hygrothermal conditions and high temperatures are required to extend the use of CFRP to high performance. In terms of mechanical properties, thermo-oxidative stability and processability, PMR-15, an end-capped bismaleimide polymerising system developed at the NASA Lewis laboratories, is being commercially developed for continuous use at temperature near 300·C under oxidative conditions /1/. To fully establish the applicability of these materials, it is essential to identify the maximum temperature at which the material can be usefully employed. In this paper, we present the results of a study of the thermal and micromechanical degradation processes under the inert and oxidative conditions.

596

2. Experimental

The laminates for thermal analysis were fabricated from prepreg made from~ woven carbon fibre in an autoclave at 288°C, followed by post-curing at 315'C in a thermally programmable oven in air. The 0' unidirectional and 0'/90'/0° cross ply laminates, on which the retained mechanical properties after isothermal ageing in air were studied, were made from unidirectional prepregs of the same materials; but they were cured at 327'C in a press-clave and post-cured at 330°C in air. All the laminates ranged in thickness from 2.00 to 2.22 mm, and contained 31 to 33 wt% of resin, which corresponds to a fibre volume fraction of approximately 60%. The glass transition temperature of the matrix increased from 350 to 370'C with a change in post-curing temperature from 315 to 330'C as determined from the maximum in Tan 8 by DMTA.

Isothermal and programmed thermogravimetric analysis on a Stanton Redcroft TG762 thermo-balance was used to identify the temperature at which rapid decomposition of the resin took place, and to monitor the rate of degradation of the laminates at specified temperatures.

For tensile testing, coupons measuring 200x20 mm2 with adhered strain gauge and aluminium end-tabs were tested in tension at room temperature, after isothermal ageing, on a Mayes, SM200 at a strain rate of 0.08 %/min.

3. Results and Discussion

3.1. Decomposition and Kinetics of Degradation

As shown in Fig.l, rapid decomposition occurs at about 420°C independent of atmosphere. At temperature below 400'C, weight losses in air and in nitrogen were identical at about 0.7 wt%, which could be attributed to trapped volatiles and/or residues from incomplete resin reaction during fabrication. Above 420'C, the rate of degradation in air was larger than in nitrogen, which is due to oxidation. As shown in Fig.2, under isothermal conditions, oxidation can be significant below 420'C. A similar measurement at the glass transition temperature (350'C), however, showed only a small weight loss of 2.3 wt% after isothermal ageing for six days, demonstrating good thermogravimetric stability of the laminates below the glass transition temperature.

Over the temperature range of 382 to 457'C, the degradation curves showed two noticeable changes in the rate of degradation with one at a weight loss of about 4 wt% and another at about 33 wt%, suggesting a three stages process: I, an initial stage up to weight loss of about 4 wt%; II, an intermediate stage at a higher rate, until the sample weight had been reduced to about 33 wt%; and III, a final stage in which the rate of degradation became slower. Within each stage the rate of degradation is approximately ~onstant. It is identified that oxidation and volatilization of the matrix was the

597

dominant degradation process both at stage I and at stage II. The apparent rapid rate of weight loss during stage II is associated with the increase in surface area created by the formation of microcracks in the laminates mainly at the interface between fibre and matrix /2/, which also make the diffusion process of oxygen much easier. Since the weight fraction of resin in the laminates is approximately 33 wt%, stage III was attributed to the fibre oxidation, which is significant even at a temperature as low as 382'C.

3.2. Thermally generated damage

3.2.1. Unidirectional laminate

Above the initial glass transition temperature of the matrix, the surface resin is found to oxidatively removed from the laminates. This is followed by the formation of extensive longitudinal microcracks mainly at the interface between the fibre and matrix, which is seen above to facilitate further degradation. Above 420'C, however, matrix decomposition causes bubbles to form inside the laminates even during the initial degradation stage. The stress-strain relationship reflects these changes so that the longitudinal tensile strength is reduced almost linearly to 72% of its original value after isothermal ageing at 457°C in air for 205 mins, whereas the tensile modulus remained unaffected /2/.

3.2.2. 0'/90'/0' cross ply laminates

Fig.3 shows a penetrant enhanced X radiography of a cross ply laminate of (02/902/02/902/0')s configuration after isothermal ageing at 400'C in air for 1 hour (corresponding to thermal degradation stage I). It is evident that microcracking took place both in the transverse direction and in the longitudinal direction. The microcracks in the longitudinal direction had already extended to the full length of the specimen, in contrast to the microcracks in the transverse direction which had been initiated from the free edges. Microscopic inspection of the polished edge of a series laminates aged individually at 400'C in air for differing periods revealed that the density of these transverse microcracks was larger in the outer 90' layers closest to the laminate surfaces. A gradual reduction in density was observed for the inner 90' plies. Fig.4 shows clearly that the microcrack density in the outer 90' layers is higher than in the inner 90' layers. These results are further supported by the nature of the transverse cracking of the outer 0' plies (Fig.3). After ageing for 12 hours, both the penetrant enhanced X radiography and optical micrography confirmed that the microcracking had become saturated throughout the thickness (Fig.5). Despite extensive microcracking, the tensile modulus was not significantly affected, but the ultimate tensile strength was reduced by about 14%.

The microcracks which are observed after isothermal ageing result from the transverse thermal residual strain which is induced by cooling the laminates to room temperature for examination, and arises from the mismatch in the thermal expansion coefficients of the 0' and

598

90' plies /3/. Therefore, the cooling interval will be determined by the stress free temperature of the laminate which is closely related to the glass transition temperature of the matrix. The Tg was found to increase gradually from 372'C to 433'C during isothermal ageing for 4 hours. Correspondingly, the stress free temperature will have increased, resulting in a larger thermal residual strain and an increased probability of transverse cracking. The increase in glass transition temperature can be understood in terms of enhanced curing. However, thermal equilibration is expected during the first hour so that the differential cracking is more likely to result from either a geometrical constrained effect, a variable cure or the presence of a plastic ising volatile. The latter would appear to be more probable since a) we observe bubble formation within the laminate /2/ and b) the curing chemistry is known to involve volatile formation and subsequent copolymerisation. Further research is required to identify these phenomena and to quantify the ageing process.

Conclusions

The thermal stability of the PMR-15 based carbon fibre bomposites has been evaluated at temperatures above the initial glass transition temperature of the matrix. Rapid decomposition of the matrix occurs at a temperature of 420'C independent of atmosphere. Under these oxidative conditions, three degradation stages have been identified. The material is essentially thermogravimetrically stable at temperatures at or below 'the matrix glass transition.

The development of damage in the unidirectional laminates at temperatures above Tg is controlled by the degradation and microcracking of the matrix. However, for the cross ply laminates, microcracking results from thermally generated residual strains which reflect thermomechanical changes in the properties of the matrix.

Acknowledgement

This research was carried out with the financial support of the SERC and the Ministry of Defence under the co-operative scheme.

References

1. T.T. Serafini in ICCM Ed W.C. Harrigan (AIME, Warrendale, USA) 1985, pp1007-1023.

2. F.R. Jones & Z. Xiang in proceedings of 3rd Int. Conf. FRC'88-Extending the limit, PRI, Liverpool UK 1988.

3. M. Mulheron, F.R. Jones & J.E. Bailey Compo Sci. and Technology 25 (1986) 119.

599

19 r--------------------------------,

15 )( In air

• In nitrogen

~ 11

" " 0 ..J .. .c .2' ., 7

~

3

-1 o 200 400 600 800

TemperatlWe ( OC)

Fig.1 Thermograms of.PMR-15 Based Carbon Fibre Composites

64

S~ge 18 48

~ • • 0 ~

i "i ~

S~geH

16

0~~~~~---_________ ~ _________ ~ ______ LJS~~~9~.~I ______ ~

o SO 100 150 200 250

TIme (hrsl

Fie.:.2 The Rates of Degradatioll of PMR-15 Based Carbon Fibre Composites At Elevaled Temperatures in Air

600

Fig,3 Microcracks in the longitudinal and transverse plies by X-radiography after isothermal ageing for 1 hour at 400 'c in air

Fig,4 Microcracks in the transverse plies by optical micrography after isothermal ageing for 1- hour at 400 "C in air

601

a) X-radiograph

b) Optical Micrograph

Fig.5 Microcracks in the cross ply laminates after isothermal ageing for 12 hours at 400 °c in air

THE HYGROMECHANICAL DEGRADATION OF ARAMID·EPOXY COMPOSITES

w. JANSSENS, I. VERPOEST, L. DOXSEE

Catholic University of Leuven Department of Metallurgy and Materials Science

De Croylaan 2 . 3300 HEVERLEE·LEUVEN . Belgium

ABSTRACT.

The presence of moisture in a composite causes the composite to expand and to lose some of its strength. In composites made with glass or carbon fibers, the moisture is absorbed by the matrix alone. In aramid-epoxy composites, the moisture is absorbed by both the epoxy matrix and the fibers. It has been proposed that a surface treatment of the aramid fiber may improve the hygromechanical properties of aramid-epoxy composites. An analysis of these properties is made and a correspondence between the location of fracture and the local moisture content has been established.

1. INTRODUCTION.

The hygromechanical behaviour of aramid-epoxy composites is inves­tigated from different points of view. First the mechanical strenghts were measured. The evaluation is based on the global moisture content of the specimens 111. Secondly, the fracture modes in the mechanical tests were detected from a microscopical investigation. Based on the calculation of the moisture concentration profile at different stages of the moisture absorption 12/, the local moisture content at the location of fracture is determined. The presentation of the strength values as a function of the local moisture content, shows the impor­tance of this approach. It turns out that in order to compare the results of tests performed at different levels of relative humidity and at different temperatures, and certainly to extrapolate the labo­ratory results to real-life-conditions, one has to know the actual moisture content at the critical locations.

604

2. MATERIAL.

The material used is TWARONR HM aramid fiber imbedded in an epo>:y system manufactured by Ciba-fieigy. The fiber content is 60 i. for all materials. An overview of some mechanical properties of the different constituants is given in table 1. In order to evaluate the influence of the interface on the moisture absorption and related aspects, both untreated and chemically treated aramid fibers are studied.

3. EXPERIMENTAL PROCEDURE.

3. 1. Temperature and relative humidity.

In this study, specimens of unidirectional composite materials were subjected to four different environmental conditions for up to four months. The conditions used consist of four combinations of tem­perature and relative humidity, namely, 76 'C combined with 70 i., 85 i. and 95 i. relative humidity and 50 'C with 95 i. relative humidity.

Two HEREAUS moisture chambers, type VLK 04/150, were used for the non-95i. conditions, while the highest relative humidity was realised in two self-designed chambers, controlled by two HEREAUS heating ele­ments and an external hygro-thermo-meter ROTRONIC Hygroskop DV-2. The specimens were measured with a METTLER AE 240 analytical balance with a standard deviation of 0.02 mg. The mechanical tests were performed at room temperature.

3.2.The interlaminar shear and transverse tensile strength.

The mechanical degradation was investigated by means of the short beam shear· test, known as ILSS test, from which the interlaminar shear strength S.y is determined 13/ :

where

( 3 * P ) * { 1 + ( P / Py ) (1)

( B * n • h )

nand h are the width and height of the specimen, respectively 6.35 and 3.00 mm P and Py are the respective loads ( in N ) at fracture and at the first deviation from linearity in the loading curve

Also the transverse three point bending test was performed, measu­ring the transverse.tensile strength S~, based on simple beam theory:

where

3 * P * 1 (2)

2 * n * h 2

nand h are the width and height of the specimen, respectively 6.35 and 2 mm

is the span-length equal 10 mm P is the maximum load ( in N )

605

Each time seven specimens were tested, so a meaningfull statis­tical analysis is possible. The global moisture content M of the spe­cimens is derived from their initial dry weight Go and the actual weight G during the moisture absorption:

M G - Go )

---------- * 100 % Go

(3)

The mechanical properties were measured in dry state, during the moisture absorption, after reaching the equilibrium level, during the recorded second moisture absorption above the equilibrium, and after redrying from a given moisture content.

3.3.The fractography.

Photographs of the fracture surfaces were taken with an lSI SS60 electron microscope and a stereo-optical microscope. Both dry and moist composites were investigated.

3.4.The moisture concentration profile.

Based on the results for both the diffusivities and the equili­brium moisture contents, the concentration profiles of the moisture through the thickness of the specimens for the mechanical tests were calculated.

4. RESULTS.

4.I.The moisture concentration profile.

It is clear that, in the initial state of the moisture absorption, the moisture distribution is highly non-uniform, but even up to 90 per cent of the equilibrium level the difference between the moisture con­tent at the edge of the specimen is still 25 per cent higher than the mid plane level, see figure 1. This turns out to be quite important in the evaluation of the results of the mechanical tests where fracture takes place in the inner part of a specimen.

4.2.The mechanical properties.

Attention should be payed to the difference between the global moisture content M, and the calculated local moisture content M, .. ". In the figures the x-a;:ls presents the moisture content while the y-axis corresponds to the relative strengths. The mid points represent the mean value of seven tests, while the upper and lower values are calcu­lated with the student-t-test, with a confidence level of 95 7., and the Q-test is used to check the rejection of the extremes with 90 % confidence. The arrows on the curves represent the redrying of the specimens from a moisture content indicated by the beginning of the arrow, to the level indicated by the arrows end.

606

4.2.1.The interlaminar shear strength.

The interlaminar shear strength S.y of the aramid epoxy composite containing treated fiber is 40 per cent higher than S.y for the compo­sites with the untreated aramid fiber. This conclusion is valid in both the dry and the moist states, see figures 2 and 3. Thus the rela­tive degradation of both composites, given by (S.y)/(S.v)O, is almost equal. The higher absolute strength of the treated aramid fiber is also reflected in the fractographical behaviour. For the untreated fiber there is an interface failure in both dry and moist state. The treated fiber splits longitudinaly when tested in the dry state and the degradation of the interface is substantially lowered during mois­ture absorption, compared to the composite containing untreated fibers

The second important difference between treated and untreated aramid fibers is that the composites which contain the latter, after redrying, regain only up to 68 per cent of their initial interlaminar shear strength, while the composites containing treated aramid fibers regain up to more than 80 7. of their initial value see figures 2 and 3. Thus the fiber treatment improves the absolute interlaminar shear strength of a redried specimen by more than 60 per cent.

The importance of interpreting the results of the mechanical tests by means of the local moisture content, is clearly demonstrated by the graphs 4 and 5 compared to respectively 6 and 7. The values of the interlaminar shear strength of the composites conditioned at the high­est relative humidity show less degradation, compared to the compo­sites conditioned at a lower relative humidity, when the overall mois­ture content is taken into account. A presentation of the results in function of the local moisture content shows perfect agreement for the different relative humidities. This demonstrates the importance of the influence of the non-uniform moisture distribution. The best way of comparing, and even extrapolating, the results of various conditions, is based on the local moisture content within the composite, instead of calculating the overall moisture content, as is done in most pre­vious studies on hygromechanical properties, encountered in literatu­re.

4.2.2.The transverse tensile strength.

The same conclusions as for the ILSS test are almost valid for the transverse tensile stength. The difference in strength for composites containing untreated aramid fiber compared to the stronger ones con­taining treated fibers is also 40 7., see figures 8 and 9.

The fractographical analysis shows the different failure modes of both composites. The naked fiber of the untreated aramid-epoxy compo­sites clearly shows interface failure in both dry and moist state, while for the treated aramid-epoxy composites there is a clear fiber splitting in dry condition, which diminishes during moisture absorp­tion.

607

5. CONCLUDING REMARKS.

One can conclude that for the evaluation of the hygro-thermo-me­chanical behaviour of a composite material, it is possible, by means of an analytical descripition of the moisture absorption, and a calcu­lation of the moisture profile, and after performing many mechanical tests, to quantify the degradation of the strength of the composite and the permanent damage after redrying. Both absorption and desorp­tion effects have to be studied. The knowledge of the actual moisture content at the location of fracture turns out to be very important, especially for extrapolation of the test results to real-life-condi­tions.

b. ACKNOWLEDGEMENTS.

The authors wish to express their gratitude to 6.S. Springer, B. Naughton and F. Elkink for the interest and the fruitfull discussions.

7. REFERENCES.

1. I. Verpoest, G.S. Springer, Moisture absorption characteristics of aramid-epoxy composites, Journal of Reinforced Plastics and Compo­sites, Vol. 7, January 1988.

2. G.S. Springer, Environmental effects on composite materials, vol. 1 Technomic Publishing Company, Inc., Westport U.S.A., 1981.

3. I. Verpoest, G.S. Springer, Effects of moisture on the compressive and interlaminar shear strengths of aramid-epoxy composites, Jour­nal of Reinforced Plastics and Composites, Vol. 7, January 1988.

60 III

! 40

>< 20 til til

6

.oJ ~ Q)

3.0

.oJ 2.0 .S u

~ 1.0 ~ .oJ til

.~ 0.0 +fl-~--~ff=~~~r-~~rh il O.

Figure:

width of the specimen (mm)

: "oisture concentration profile Untreated ar u~i d-epoxy (oliPosi te

o. "tion: 50 'e * 95 f. R.H.

I 60

~ III f f

. y

i ! 40 0

~ >< til 20 Ul

0 0.0 2.0 4.0 0.0 2.0 4.0

Figure: 2 : Moisture content M

Interlalinar shear strength Ssy versus overall loisture content H Untreated aralid-epoxl cOlposi te Condition : 76 'C • '1J I R. H.

Figure: Moisture content M

.) : Interialinar shear strength Ssy versus over all loi sture content Treated aralid-epoxy cOlposite

n .. • • c

608

LO_ t 1 I 0 !

>. ~. til 0.5-til

t ~ "->. til 0.00 e .V "'~V 4'.V

Moisture content M rigl1re : 4 : !nteriar.inar shear strength Ssy/Ssyo

versus everali loisture content 1\ ynt~~!~ed ar~~l ~:epoxl ~o~po5i te I onu; clon : ,6 l I 9J • R.H.

1.0 _ r 1

of: .~0.5-~, >; , l2 0.0 - u.u "'.V '1.V

0

~ til -"->; til e O.OO~--~~~--__ ~~~~ .U "'~U 4;U

Moisture contant M fIgure: 5 : Interialinar shear strenoth SsylSSYo

versus overall loisture 'content II ~ntr~.ted ara~id-epoxy coap,osl te Londlhon : 76 'C i 70 1 R.H.

o

~ til ... "-~ til 0.0 .,n-,rr-~-rTl~~-,.......r--l - v.v "'.V 4.U

Local Moisture content Mloc Local Moisture content Mloc FiQu'< : 6 : lnterla!inar shear strength Ssy/Ssyo

versus iocal loisture content "'ac Untreated ar.lid-epoxv composite Cendition : 76 'C i 95 l R.H.

40-

~ 20-

o u~.u~~-~~.~'u--~-~~ .• ft-u~

Moisture content M Figure: a : Traflsverse tensile strenQth St

versus overall loisture content K tint~;a~ed ~r~~i~~epc~t coeposite Con itiDn . 16 _ f 9) 4 R.~,

Figure: 7 : Interlalinar shear strength Ssy/Ssy versus local lUi sture content H, ac Untreated aralid-epoxy cOIposite Condition: 76 'C • 70 Z R H.

40 ..... I -;0 f

~ 20_ if •

.jJ til

0 O.U ~.U 4.'U

Moisture content M Figure: 9 : Transverse tensile strength St

versus overall loisture content iI Treated aralid-epoxy cD.posite Condition: 7b '[ l 95 1 iI.H.

Table 1 : Properties of the fibers iod the epoxy resin

Density IIodulus Tensile Fiil .. e Str~h Strain

(g/c~1 (6Pal IlIPil III

Resin ~ox~ b~ Cibi 6ei3r ( Y5 6/ Y917/DY07 1.15 3.4 8:i 4 ( 100 I '10 I 1 I

Fibers Araaid TMARONR HI! 1.45 125 3700 2.7 by EI«A-AtCZO

FATIGUE BEHAVIOUR OF GFRP: SOME CONSIDERATIONS ABOUT INTERFACES

L. VINCENT, L. FIORE, P. FOURNIER

Ecole Centrale de Lyon BP 163 - 36 Avenue Guy de Collongues - 69131 ECULLY - France

ABSTRACT

New considerations about fatigue behaviour of GFRP through the roles of the

interfaces are presented. For bending fatigue tests, interfaces are described as

privileged ways of cracking. Developed from the main idea of the "resistance of the

materials " to the cracks nucleated at the very first loadings, this work analyses the

subsequent growth of cracks and damage : weaker shear strengthes in interfaces can

induce higher lifetimes.

I - INTRODUCTION

The fatigue behaviour of composite materials is often related to the

fiber-matrix interfaces. Three years ago, we showed that the rate of the coupling

agent can strongly modify the fatigue curves and that an optimum rate exists

(P. Jeanne et coiL) in the case of unidirectional R glass/epoxy composites.

Unfortunately, the role of the interfaces is still unknown and it is very difficult

to correlate physico-chemical properties as defined on model products and

macroscopic mechanical measurements : interfaces are well known to transfer the

applied load if we consider the global strength but their resistance to the early formed

cracks is too often neglicted. The role of the interfaces is discussed following three

considerations about the fatigue behaviour of UD (0°) GFRP:

610

1 - the fatigue behaviour of GFRP must be analyzed in terms of " Tolerance to

the first nucleated cracks"

2 - the analogy between the failures appearing under a monotone loading

("static" failure) and the fatigue tests .

3 - fatigue damage leads to break away in the global behaviour the

contribution of each constitutive element.

2 - FATIGUE BEHAVIOUR OF GFRP : THE ROLES OF THE INTERFACES

2.1. Tolerance to the first nucleated cracks

Keeping in mind that fiber and matrix breaking can occur on the stretchest

surface of the specimen as soon as the first loading (P. Puget et coli.), we identified

three main mechanisms of failure: fiber breaking (filament failure ie 10 to 20 IJ.ITl in

length); matrix cracking (below one millimeter in length) and superficial debonding

parallel to the fiber reinforcement (one or two millimeters). After this first stage,

we noticed two extreme developments : multiplication of the failures or localized

crack growth:

o ) Tranfert loading and accommodation of local strain through the interfaces

induce a multiplication of the numbers of broken fibers :

micrometers in length debonding can accommodate the local strain

and thus increase the lifetime.

at the opposite, if no debonding can develop, no local accommodation

is possible: then the first crack grows from these early formed

defects.

o ) Localized crack growth in the matrix leads to a higher environment

effect (moisture) both in fiber (delayed fracture) and matrix (plasticization) .

Here the role of the interfaces as a chemical protection of the fibers becomes

more important.

In . bending we must emphasize that privileged ways of cracking in

interfaces can Induce a large Incease in fatigue lifetime and strength.

611

2.2. Analogy between the failure under the monotone loading ("static" failure)

and the fatigue test.

Here we mainly insist upon temperature and strain rate (or frequency for

fatigue). We find an increase in the ultimate bending stress as temperature decreases

or strain rate increases (the apparent modulus remains constant) (fig.1). Once more,

mechanical values can be related to the failure features. In both cases, the highest

stresses are associated to feather-like features.

These results well agree with the well known viscoelastic properties of

polymeric materials according to temperature and frequency dependences. They

strongly suggest that :

the viscoelastic response of the epoxy matrix plays a major part in the

composite behaviour,

the interfaces can be privileged crack paths.

Bending fatigue tests (R ratio-that is tmin/tmax- is chosen between 0.1 and

1) indicate poorer behaviours with decreasing test temperatures down to - 60°C (fig

2). Tendencies are not clear enough for the upper room temperatures (up to 60°C) :

results are in the scattering bands of room temperature tests (23°C). Besides, we

must notice that the rates of damage (as measured by stiffness reduction) are quite

different : about 10-6 %/cycle for 23°C up to 10-5 %/cycle for - 60°C.

Tests run under different frequencies (10-25 Hz) show the same tendencies in

SoN curves (criterion of a 10 % drop in stiffness). As for tensile tests ( G.D Sims et

coli). we must define the fatigue level through the ultimate strength obtained under

the same strain rate or temperature. This appears very important specially for

GFRP where the rate of .loading increases the fiber strength: about 40 % from 0.2 to

200 MPa/s (J.E. Ritter et coli.).

Here the time above the critical strain allowing the propagation of

defects is more significant (analogy is to be made with static fatigue -delayed

fracture -of ceramics)

Any explanation of the fatigue behaviour needs also to take into account the

features of the first cracks appearing on the stretchest surfaces at the first cycles.

612

According to strain rate and temperature conditions, viscoelastic properties of matrix

and interfaces are modified during bending fatigue test and thus contribute to stiffness

reduction with various damaging kinetics.

The physico-chemical role of the interfaces is also pointed out through static

fatigue tests (R = 1) under several humidity conditions.

For the highest level (RH >90%), the drop in stiffness is much greater than for the

low RH percentages (down to 0%) at the same room temperature (fig.3). Recent

results confirm the effect of water on damage kinetics specially when cracks on the

stretchest surface remain open (R=1). This effect can be related to a delayed fracture

mechanism of the glass fibers (S.M. Wiederhorn) and also to a plasticization

(increase in ductility of the matrix) in front of crack : for instance cracking in the

interfaces can occur due to the imposed local strain or due to the fretting wear

between fiber and matrix (cyclic test). Here, the global increase in ductility of the

matrix is not quick enough to justify such stiffness drop (below few hours).

These experiments pOinted out a main parameter (humidity) in

predicting fatigue behaviour and durability of industrials parts.

2.3. Fatigue damage leads to break away in the global behaviour the

contribution of each constitutive element.

The behaviour of both matrix and interfaces is also analysed through damping

capacity measurements (in situ or delayed testing), we notice that:

-The restored and dissipated energies (in situ measurement of the 0'-£ cycles

by the means of the Rheometer(L.Fiore)) can quantify "the number of cycles" during

which matrix and interfaces are main elements of the material response. For instance

in pultruded E glass/epoxy (60%v/o), we show a shuttering of hysteresis loop during

the fatigue process (fig.4). This results are well correlated with WOHLER's or

MANSON-COFFIN's curves in which the contribution of the matrix seems very weak

(L. Fiore).

- The damping measured through dynamical viscoelastic analysis (delayed

testing-after several levels of fatigue damaging- by the means of PL-OMTA (B.

CHABERT and J. CHAUCHARO's laboratory in UCB-L YON(F)) indicates the "integrity"

of the material. Pultruded GFRP (60%v/o) are characterized by a weakest

participation of the fibers to the reinforcement.

613

The global viscoelastic behaviour depends on the increasing number of broken fibers:

with cycles material becomes less and less "bound" (fig. 5)

Damping analysis (by the Rheometer cycle or the DMTA) give us

complementary informations both related to the breaking away in the global

behaviour of the contribution of each constitutive element.

This approach can favour a more rational analysis of the nature of the resins or

the interfaces that can improve the fatigue behaviour.

Static and fatigue (R = 0.1 to 1) bending tests have been run on several UD

glass-epoxy composites with an accurate analysis of the initial damage mechanisms.

The interfaces appear as a main element in the transfer of load between the fibers but

also as privileged crack paths in a "Safe Crack Growth-like" approach.

The interface role has been studied using several test conditions: temperature,

strain rate and humidity. The first two parameters indicate the influence of the global

mechanical response of both matrix and interfaces. The latter shows the chemical

effect of the interfaces and their possible protective role towards the fibers.

These considerations (damping and stiffness reductions versus test times,

time-temperature dependence) must be taken into account for any fatigue model

described from a continuous damage laws. Particularly the fact that the interfaces can

be privileged crack paths is very important to be included in mean life predictions.

REFEREflCES

1 - P. JEANNE et ai, ECCM 1, Bordeaux (1985) 158-163.

2 - L. FIORE, thesis 88-20, Ecole Centrale de LYON (1988)

3 - P. PUGET, L. FIORE, L. VINCENT, JNC5, Paris (1986), 715-728.

4 - G.D. SIMS and D.G. GLADMAN, NPL report DMA (A) 59 (1982) 1 -24.

5 - J.E. RITTER, J.M. SULLIVAN, K. JAKUS, J. of Applied Physics, vol49

(1978) 47-79.

6 - S.M. WIEDERHORN, Fracture mechanics of ceramics, vol2 (1974) 613.

614

T:23·C

l/W 20

~ _______ ~0~.5 ______ ~1 V - MM.'S

800

4CO

<rMP' A

V2MMlMN

•

Fig.1-Loading rate and temperature effects in bending static test for UD E glass/epoxy

100 FiFo %

50 ~:22~ V.25HZ R:o.1

lll+20

Fig.2 - Temperature effect

on bending fatigue (R=O.1) FdaN 1700c 191 ___________ _

2. mm

Fig. - Load (F)-displacement (f)

shape versus cycles

N

119

0.05

f. :2.5%

4-i:20

50 100 th

Fig.3-Humidity effect

on static fatigue(R=1)

Fig.5- tga for three

stiffness drops

SYSTEMATIC FRETTING WEAR AND FRETTING FATIGUE STUDIES ON CARBON FIBRE / EPOXY LAMINATES

O. JACOBS, K. FRIEDRICH, K. SCHUL TEo

Technische UniversiUit Hamburg-Harburg Harburger Schlossstrasse 20 - 2100 HAMBURG 90 - West Germany

"DFVLR - Linder Hohe - 5000 KOLN 90 - West Germany

ABSTRACT

Fretting wear and fretting fatigue often limit the lifetime of structural parts under cyclic loading. Although fibre reinforced plastics are widespread in such applications, there is only little known about their fretting wear and fretting fatigue performance. The present paper investigates the influence of an additional fretting component on the fatigue behavior of a CF/EP -laminate. A quantitative model for the degree of fretting fatigue damage is proposed. Separate studies on the fretting wear of these materials are employed to explain the fretting fatigue results.

INTRODUCTION

The expression "fretting fatigue" describes the following loading situation: a fatigue loaded structural part makes a cyclic strain motion which produces an oscillatory sliding in the contact region to a counterpart. The surface damage induced may penetrate into the bulk material, forced by the fatigue load, and it may result in sudden failure. For this reason fretting fatigue is of great importance in practice (e.g. multi layer leaf springs, bolted jOints. flanges) and attracted extended research activities in the case of metals 11/.

A pilot study about the fretting fatigue performance of CF/EP - laminates indicated that the fatigue life of these materials may decrease if an additional fretting component damages load bearing 0°-layers 12-4/. The current project intends to investigate systematically the interaction of fretting surface damage and material fatigue. The influence of contact pressure and counterpart material is studied.

616

I - EXPERIMENTAL

1.1. Materials

For the fretting wear tests a simple unidirectional laminate (L 1) was chosen. The fretting fatigue experiments were performed with a cross ply laminate [02,902125 (L2) and a quasi isotropic laminate [02,±45,02,±45,9013s (L3), both of them with covering 0°-layers. All laminates contained T300 carbon fibres. For the laminates L 1 and L2 the epoxy resin R 5212 (BASF) was used as matrix, whereas L3 possessed a Fiberdux 914 C (Ciba Geigy) epoxy resin matrix.

The fretting pads were cylindrical pins with a diameter of 5 mm and consisted of aluminium (Vickers hardness Hv = 103) or stainless steel (Hv = 305). The front sides were ground and polished.

1.2. Testing procedures

A SRV testing system, commercially available by Optimol Instruments, was employed to carry out fretting wear tests. The laminates were cut to disc shaped specimens which were fixed on a specimen holder. Fretting pins slid oscillatory with their front sides against the specimen surfaces parallel to the fibre orientation. The full oscillation width was 500 J.Lm.

For the fretting fatigue tests the laminates were cut to beam shaped specimens with a width of 8.3 mm (L2) and 6.3 mm (L3), respectively. Fretting pads, hold by a special device at a fixed position, were pressed against the fatigue loaded specimens 12/. A fatigue load was applied in the tensile region with a frequency of 10 s-1. The ratio of lower to upper fatigue load was set to R = 0.1. The specimens strain amplitude in the fretting region was measured to be about 700 J.Lm.

" - RESULTS AND DISCUSSION

2.1. Fretting wear

Figure 1 shows the mass loss of a CF/EP - laminate as a function of time and counterpart material. Two distinct phases of the wear process occur. In a rather short "running in period" 151 the material is worn very rapidly. After a certain time a steady state is reached and the wear process proceeds significantly slower and linearly with the time. Obviously the wear of the samples is higher for the softer counterpart material. The reason consists in the roughening of the pin surface by the abrasive fibre debris. This process acts more effectively the softer the pin material.

The contact pressure between the fretting pin and the specimen has a critical influence on the fretting wear. Below a contact pressure of about 20 MPa the specific wear rate 151 remains constant at a rather low level but increases above this limit suddenly by a factor of about eight (fig. 2). Such

617

limiting p-values are well known in wear processes 151 and depend in their magnitude on the sliding speed and environmental temperature.

2.2. Fretting fatigue

Figure 3 presents a plot of the applied upper fatigue load versus the resulting lifetime (WOhler-curve). This diagram leads to the following conclusions:

1.) At contact loads up to 450 N (22.9 MPa) the use of stainless steel pins leads to fretting fatigue lifetimes in the scatter band of the pure fatigue life. This corresponds with the high fretting wear resistance of CF/EP versus stainless steel.

2.) In the case of aluminium pins no significant fatigue life reduction occurs below a contact load of 200 N (10 MPa). But between 200 Nand 360N (18.3 MPa) the lifetime drops by some orders of magnitude.

3.) To find a quantitative description of the fretting fatigue damage one can use a relatively simple model. In first approximation the deviation of the fretting fatigue from the pure fatigue curves should depend on the damage of the load bearing OO-Iayers 13/. During the test the two fretting pins penetrate into the covering layers and thus reduce the total cross section Ao of the 00 -

layers by an amount 2 M. As a result the upper fatigue load at which the specimen fails after a given number of cycles decreases by the same proportion as the effective cross section of the OO-Iayers in the fretting region. This leads to the definition of a "relative fatigue strength reduction":

A ~ _ 0F- OFF _ 2 Il A LlVrel- - *-

OF Ao OF = upper fatigue load at which the specimen fails after N load cycles 0FF= upper fretting fatigue load at which the specimen fails after N cycles

(1 )

Figure 4 shows Ilorel as a function of the fretting fatigue life N. All curves start in the origin and increase proportionally with testing time. In contrast to pure fretting wear no running in period occurs. From the slope of this linear part of the curve some kind of "specific pseudo wear rate" can be derived with equation 2:

0* 1t*IlOrel*Ao*d W - -----=---

s - 16 F N * S • IlN (2)

d = pin diameter, FN = normal load, S = full oscillation width .. 700 11m.

Figure 5 presents the results of these calculation. Like the specific fretting wear rate the specific fretting fatigue pseudo wear rate increases stepwise at a critical contact load. But for fretting fatigue this critical load is

618

reached already at about 15 MPa, Le. at a clearly lower magnitude than in the case of pure fretting. In addition the step in Ws is significantly higher for the fretting fatigue situation. Probably the fatigue load amplifies the fretting surface damage and supports its penetration into the bulk material.

The l\crrel - N - curve reaches a plateau, after the two fretting pins have penetrated totally through the four covering OO-Iayers. The fretting damage of the following 90 0 - or 45°-layers respectively affects the laminate strength insignificantly /3/. In this case the saturation value should amount to about

l\crrel = (d/w)*(npnT) = (5/8.3)*(4/8) = 30 % (3) d = pin diameter, w = width of the specimen, nF = number of covering 00 -

layers, nT = total number of OO-Iayers. This value is in good accordance with the experiments. Figure 6 presents a l\crrel - curve for laminate L3 which is derived in the same manner as the curves in figure 4. This curve shows qualitatively the same course as the AI 400 N - curve in figure 4, but since laminate L3 possesses 24 OO-Iayers and the specimens are smaller, the plateau should be reached at

l\crrel = (5/6.3)*(4/24) = 13.2 % (4) This value again is in good agreement with the experimental results.

III - CONCLUSION

In agreement with predictions /3/, the fatigue life of CF/EP-Iaminates is reduced due to an additional fretting component. The occurrence and magnitude of the fatigue life reduction is correlated with the fretting wear behavior of the samples. Low fretting wear leads to low or even insignificant lifetime reductions; parameters such as pin material or contact load which affect the fretting wear act critically on the fretting fatigue performance too. However, the fretting fatigue damage follows a different dynamic than the fretting wear and proceeds much faster. As a good quantitative measure of the fretting fatigue damage the "relative fatigue,strength reduction" was proposed.

IV - REFERENCES

1. R.B.waterhouse (ed.): Fretting Fatigue, Appl.ScLPubl., London, 1981. 2. K.Friedrich, S.Kutter, K.Schulte, Compo ScL Technol. 30 (1987) 19. 3. K.Schulte, K.Friedrich, S.Kutter, Compo Sci. Technol. 30 (1987) 203. 4. K.Schulte, K.Friedrich, S.Kutter, Compo Sci. Technol. 33 (1988) 155. 5. K.Friedrich, O.Jacobs, in Verbundwerk 88 (S.Schnabel, R.Gadow,

J.Kriegsmann, ed.), 1988, VERBUNDWERK, Wiesbaden (FRG)

ACKNOWLEDGEMENTS

The project is financed by the Bundesministerium fOr Forschung und Technologie (FRG), project no. 03 M 1022 A. Some of the materials are kindly supplied by the BASF.

t 1,000

r;; 0,800

oS 0,600 UI ~ 0,400 -I UI 0,200 UI

~ 0,000

~ ..

Va ~

""'" - .. o 2

619

-- a Steel Pin .. Alu Pin

4 6 8 10 12 Time [h] •

Fig.1 - Mass loss of a CF/EP-Iaminate as a function of time and counterpart­material. Contact pressure p=12.7 MPa, frequency v=10 s-1.

8

E z 6 -. CO)

E E 4

'r: 5:! 2 ..... • • 11: o

Fretting Wear • Alu Pin vs. i-Laminate L1 /. •

• /

·"""i o 10 20 30 40

Contact Pressure [MPa]

Fig.2 - Specific wear rate of L 1 under fretting conditions (v=10 s-1 )

I&.

1&..650 I&.

t)

~~-L ra

1-, ~+ a Fatigue

• Alu 400 N .. Alu 300 N + Alu 360 N " Alu 200 N • 5te81450 N

~ lE!....a~ ~ ..... I--t-;-.. l .+

~., ~ .

550 10 2 1d

Lifetime [Number of Cycles]1---.... ~

Fig,3 - Wohler curves for L2 under fatigue and fretting fatigue conditions,

620

f 30 • Alu 360 N

20t-t----r-----I----~~~~~ .... ~ - 10~------r-----~~~~~-------1 ~

b <l

100000 200000 Lifetime [Number of Cycles)t--.....

Fig.4 - Relative fatigue strength reduction of L2 as a function of time and contact load

.... 60

i 50 ...... ME 40 E 30

II) o 20 ~ 10

II • J 0

Fretting Fatigue .-. Alu Pin vs. I Laminate L2

/ V -.

o 2 4 6 8 10 12 14 16 1 8 20 22 Contact Pressure [MPa]

Fig.5 - Pseudo-specific wear rate of L2 under fretting fatigue conditions.

14

12

~ 10 o - 8

CD

t5' 6 <l 4

2 o

~ J T ,

II

o

•

Fretting Fatigue Laminate L3 Aluminium Pins FN =450N

100000 200000 300000 Lifetime (Number of Cycles]

400000

Fig. 6 - Relative fatigue strength reduction of L3 versus testing time.

INFLUENCE OF MOISTURE ON THE COMPRESSION BEHAVIOUR OF COMPOSITES

ABSTRACT

G.ZIEGMANN

AKZO Fibres and Polymer Division Institute Fibres and Advanced Materials - Kasinostrasse

5600 WUPPERTAL - West Germany

In dimensioning components made with composites one has to consider, besides the mechanical stress, the influence of the surrounding atmosphere. Moisture picked up by the polymer matrix system of composites diminishes the stiffness and therefore the thermal stability of the system. The moisture absorption as well as the influence of the moisture in the laminate on the physical and mechanical behaviour are discussed.

INTRODUCTION

Fibre composite parts based on carbon fibres and polymer matrices are increasingly galnlng importance in a large nwnber of applications, in particular in the aerospace sector.

Apart from the mechanical stress to which the components are subjected, there are additional factors influencing the material in actual service. These factors, which are related to the conditions of use and the surrounding atmosphere, nearly always include the ambient moisture.

622

Moisture picked up by the polymer matrix system, however, invol ves a change in the mechanical and physical properties, which may restrict the applications for fibre composites /1-5/.

1. Materials used

1.1. Matrix

OWing to their molecular structure, polymer matrix systems have a lower thermal stability than the reinforcing fibres. This investigation cove'rs a typical example of the group of epoxy resins currently used in a wide variety of aereospace applications as well as of the BMI resins, which are able to withstand higher temperatures.

1. 2. Fibre

The experiments were confined to carbon fibres as these are the predominant fibre material used in the aerospace industry.

The specific structure of this fibre /6/ prevents the absorption of ambient moisture. So the fibres may be regarded as inert as far as the effects of moisture pick-up are concerned /5/. As described in /6,71, there exists a whole family of fibres with different tensile moduli and strength levels. The investigations reported below covered the fibre types Tenax HTA and Tenax ST3 (produced by the fibres and polymers division of Akzo).

1. 3. Laminates

The moisture uptake was determined on laminates having a multidirectional structure as described in the studies of /8/; the physical and mechanical tests were conducted on unidirectional laminates, where all fibres are arranged in one direction.

The laminates were cured according to the instructions of the prepreggers /9/. Subsequent non-destructive X-ray and ultrasonic tests revealed faultless laminate specimens.

623

2. Effects of moisture on physical properties

The mechanism of moisture absorption by polymer materials has repeatedly been described /e.g. 1-5/ and is well known. According to Fick's laws, the moisture uptake b~ such systems, when plotted as a function of time (~t-scale), shows a typical pattern as represented in Fig. 1.

Here it can be seen that the BMI system exhibits a higher rate of diffusion as well as a higher overall moisture uptake. The moisture picked up by the polymer matrix acts like a plasticizer in the rigid molecular structure; it leads to a reduction of the softening temperature and consequently to a lower maximum temperature at which the material can be used. This influence of moisture, which can be demonstrated, for instance, by dynamic mechanical analysis (torsion pendulum, DMA), is discussed in detail in /5/.

3. Effects of moisture and temperature under compressive stress

The plasticizing effect makes itself felt in particular when the material is subjected to compressive stress as the softening of the matrix may result in early failure by so-called microbuckling.

Fig. 2 shows the compressive strength and the compressive modulus as a function of temperature for the two carbon fibres with the epoxy resin system. As the temperature increases, the strength rapidly declines, while the modulus remains largely constant. This marked decline is clearly attributable to the softening of the matrix, which in the temperature range shown loses much of its supporting effect.

The compressive modulus is determined as secant modulus according to DIN 29971 and therefore is not influenced significantly by the decrease in strength.

The curves obtained for the bismaleinimide system (Fig. 3) are similar to those discussed for the epoxy resin system, although the strength and modulus decrease more slowly, which was to be expected in view of the molecular structure of the BMI system. The strong decline to a level of 50 % of the original strength at 120°C still comes as a surprise as investigations into comparable systems have shown significantly higher values /10/.

624

In Figure 4, the compressive strength has been plotted against the temperature, in relation to the starting value at room temperature, for both resin systems and fibre types. This graph, too, shows the markedly higher deflection temperature of the BMI system. Moreover, it demonstrates that the differences between Tenax HTA and Tenax ST3 found in tensile tests do no longer make themselves felt when the material is subj ected to compressive loads. The curves come very close to each other and show a similar shape.

The serious effects of moisture on compressive strength can obviously not be determined by torsion pendulum or DMA tests. As illustrated by Fig. 5, the decline in compressive strength in relation to temperature is more serious than that of the modulus in shear; a linear relationship obviously does not exist.

The compressive strength rather seems to be largely determined by the elongation of the matrix system /10/ (see Fig. 6); while the tensile strength increases with the ultimate elongation, the compressive strength begins to drop when a certain elongation level has been reached. This correlation applies only to the tests shown in Fig. 6 /10/. Whether or not it is generally valid has to be found out by further tests on other fibre/matrix systems.

4. References

/1/ R. Delasi, J.B. Whiteside: Effects of Moisture on Epoxy Resins and Composites; ASTM STP 685

/2/ A. Davis, D. Sims: Weathering of Polymers; Applied Science Publishers London, N.Y. 1983

/3/ H.W. Gitschner: Diffusionsbedingte Verformungs­und Spannungszustande in glasfaserverstarkten Verbundwerkstoffen; Dissertation an der RWTH Aachen, 1980

/4/ J. Brandt, J. Warnecke: Die Anwendung thermome­chanischer Prlifverfahren bei der Untersuchung des Feuchteeinflusses auf CFK; Vortrag zum DGLR-Symposium "Entwicklung und Anwendung von CFK-Strukturen" 08., 09.11.1984

~ ~ ~

.3 c

.2

::5

/5/ G. Ziegrnann: EinfluB der Feuchtigkeit auf den Erweichungsbereich von Epoxidharzeni Kunst­stoffe 74, 1984, H. 12, S. 732 - 735

/6/ H. Blumberg: Kohlenstoff- und Aramidfaserni Kunststoffe 77, 1987, H. 10, S. 1100 - 1104

/7 / H. Blumberg: Status quo der Entwicklung von polymeren Faserverbundwerkstoffen in der Bun­desrepublik Deutschland im internationalen Vergleichi BMFT-Tagung Sept. '88 in Hamm

/8/ J. Sauter: Studienarbeit, TU Stuttgart 1984

/9/ Chr. Seyffert: Studienarbeit, TU Stuttgart 1987

/10/ H.J. Semrau: Neue nichtmetallische Strukturver­bundwerkstoffe im militarischen Flugzeugbaui Jahrestagung der DGLR Hamburg im Okt. 1984

'.6

--- -.- ~I Ot,"'9 c:orod.ol.(lfu. 'O·C: / '~ "t. "f

625

,or" .........

/ ,-a,..1n ... [pIIIl"

• 8d_.,-.t

1II .. ~t: DftI",l £0 '11'01 '"

Fig. 1 - Moisture uptake by carbonfibre laminates with

/' ,""

/ l / II I /

Ql 'j

o o ro w ~ ~ ~ 00 ~

SI04'"lng 11m.. (rr.)

2000

'GOO i~ ~

llOO • 1'. '

~ Q"i--'ffi~;t:-:ffl.H 1 Q*--Wl~~+H.H '! QH/--I\\-'!.--':;::-'l-iJ.

1

1

different matrix systems

BOO e.g' E~

LLLD;.·~:

I

Fig. 2 - Compression behaviour of EP laminates as a function of tempera­ture

8i. '00

20 '0 00 BO 100 ,:10 "0 100 1 BO

RfSln . [PO.- ICI

rlbt .. ! T"no.ll HTA-' ~ i pno.l 5T oJ

lpslt.mp.trolur. ( 'C)

C::hmo, .. : ?o'e I 9S "I. R.F. Comprnsloo 5trf'nQth

_ __ Comprnsion modulU5

626

~ 2000 x

1600 c .

0 1200

~ ~ = !CO

~~ E ~ I.(X) 0-u"

b 1 ~ .-;}.-l.... ,g".-*

-- ' -~-~!-I ~~.

~ ~ ro ~ ~ ~ ~ ~ ~

ips.ttpm~,atutP t"C)

RUin : Bisma.llt i n ~rrUd IBM!} CI_maw- : 701: 195 -4 R,F.

' 0

Fibr. I Tt'nalll HTA-7 -- COmpfulion Itrt'f"llg1n

'i e ~ ~

'5

j G

: T'tncu ST-3 --- Cornprf'ss ion modulus

;: - - 8MI .... ~·lrt1Q,o ~-,

b 100 l--ho.,.j=4-+ - +-4 =::~::~ :: :~., J .... U' ·' ..... n. 'ltrotYS ... '

~J--+-t--': L

go 60 1-----1f-----t-4- ..., t

,, 1i .. c ~r----'I--f--+-+-+= ~ -~ ~~ 20r-----1I--f-4-+--+-+---+--f-------1--~ H '"'"

.0

30

20

10

~ ~ 60 ~ 100 120 T.mptrotur. I "C 1

= 1500 ~ CC'JrT1)tfSS ion

~ strf'nglh "MJ:

l~n Ruin : Epoxid SOO u .... fibril' . T"no:!! HTA· '

o 60

ISO

65"J.

Fig. 3 - Compression behaviour of BMI laminates as a function of temperature

Fig. 4 - Standardized compressive stress of EP and BMI laminates as a function of temperature

Fig. S - Relationship between modulus in shear and compressive strength in EP laminates

Fig. 6 - Effects of resin elon­gation on tensile and compressive properties of carbon fibre laminates

ABSTRACT

MOISTURE DIFFUSION INTO TWO-PHASE MATRIX RESINS FOR FIBRE COMPOSITES

F. JONES, P. JACOBS

University of Sheffield - School of Materials Northumberland Road - S10 2TZ SHEFFIELD - England

Moisture absorption in resins and composite materials has been an area of concern for a number of years. In most previous investigations a simple concentration independent Fickian diffusion process has been invoked to describe the moisture absorption into what is generally assumed to be a single-phase material. In this paper a model is developed which enables the diffusion characteristics of a two-phase polymeric material to be described, and the volume fraction of the phases to be calculated.

INTRODUCTION

Several authors/I-5/ have monitored moisture absorption in resins and composite materials. For simplicity most of them have assumed that moisture absorption is a simple concentration independent Fickian diffusion process. However, it is often found that the final stages of moisture absorption deviate from Fickian behaviour. Certain resin systems are known to exhibit a two-phase structure, comprising areas of differing density, which may influence the diffusion characteristics of the system as a whole. A simple model to describe the diffusion process which occurs in a two-phase resin system has been developed from studies on an isophthalic unsaturated polyester resin system, which is believed to have a two-phase structure. The validity of the model for describing the diffusion behaviour of a two­phase polymeric material was confirmed using an (ethylene,butylene)/ styrene block copolymer thermoplastic elastomer.

628

EXPERIMENTAL

Crystic 272 (Scott Bader and Co. Ltd.), an isophthalic unsaturated polyester resin, was cured with 2 parts per hundred of resin (phr) of a 50% methyl ethyl ketone peroxide solution (Catalyst M, Scott Bader and Co. Ltd.) and a 0.45phr of a cobalt naphthenate solution (Accelerator E, Scott Bader and Co. Ltd.). This formulation has a gelation time of 6-12 hours, depending on the ambient temperature, with no exotherm enabling the production of reproducible 5mm thick void-free castings between vertical glass plates lined with Melinex release film. Specimens 20x80 mm were cut from the plates using a water cooled diamond saw. All the faces were polished to a better than 1 pm finish prior to complete post-curing in an air circulating oven at 130·C for 1.5 hours. The specimens were dried in a vacuum oven at 50·C until they reached a constant weight. This procedure is necessary to ensure that the specimens are completely dry prior to testing. The thermoplastic elastomer Kraton G-1657X (Shell Chemicals U.K. Ltd.) was the two-phase polymeric material selected to evaluate the validity of the diffusion model developed from absorption studies on Crystic 272 polyester resin. Thin sheets of Kraton G-1657X, suitable for accurate absorption measurements, were supplied by Evode Plastics Ltd.. Specimens 60x20 mm were cut from the 2.73 mm thick sheets.

The specimens were exposed to relative humidities of 96% and 75%, created by the use of saturated salt solutions, and some were immersed in distilled water. The temperature during conditioning was maintained at a constant 50±1·C in an air circulating oven. The specimens were periodically weighed to monitor the absorption of moisture with time to equilibrium, (Mm). The thickness of both the Crystic 272 and Kraton G-1657X specimens when dry and at Mm were measured, and in both the systems the amount of swelling due to moisture absorption was found to be negligible.

RESULTS AND DISCUSSION

Shen and Springer/1/ have developed expressions for the diffusion of moisture in fibre reinforced composite materials which are based on the similarities between thermal conductivity and moisture diffusion in composite materials. Adopting this approach expressions for the diffusion of moisture in an isotropic two-phase system can be established. Behrens/6/ derived a general expression for thermal conductivity in a two-component system having orthorhombic symmetry. Using the thermal conductivity analogy, for spheres in a cubic lattice the expression for moisture diffusion becomes:

Ox = 01 [_( B_D_+ _2_) _+_2_{ B_D_-_l_) V_] (BD + 2) - (BD _ l)Vd

d (1)

where

629

BD = Dd/DI' D is the diffusion coefficients for the dense (d), less dense (1) and matrix (x) phases and Vd the volume fraction of' the dense phase.

The moisture absorption curve for Kraton G-1657X, in the form of 1

M(t) against (time)2/thickness, demonstrates the characteristic shape for a two-phase polymeric material (Fig. 1). Two important regions can be identified: Region I is attributed to absorption of moisture by both the dense and less dense phases of the material. In region II the diffusion process is considerably slower, and this is attributed to absorption by the dense phase alone. Assuming that initially both phases absorb moisture, but once the less dense phase has reached equilibrium the increase in moisture content of the resin results from absorption by the dense phase alone, then by subtracting the slope in region II (md) from that of region I (mx) the slope of the curve corresponding to the absorption of moisture into the less dense phase (ml) may be obtained (i.e. mx - md = ml)' Using a Fickian approach, diffusion coefficients for the dense and less dense phases (Dd and Dl respectively) can be calculated, as can the diffusion coefficient of the matrix as a whole (Dx ):

2

Dx = It mx ( 2h )]

4MID

(2 ) (3)

The thermoplastic elastomer Kraton G-1657X (Shell Chemicals U.K. Ltd.) was the two-phase polymeric material selected to evaluate the validity of the diff1ision model developed from absorption studies on Crystic 272 polyester resin. Thin sheets of Kraton G-1657X, suitable for accurate absorption measurements, were supplied by Evode Plastics Ltd •. Specimens 60x20 mm were cut from the 2.73 mm thick sheets.

and

where

(4)

MID and Ml are the equilibrium moisture contents of the matrix and the less dense phase respectively and h is the thickness of specimen.

The "edge effect" correction calculated by Shen and Springer/1/ must then be applied in order that the true diffusion coefficients may be determined. This states:

Da = D(l + (h/l) +(h/n))2 (5 )

630

where Da is the diffusion coefficient of the material in the direction normal to the plate surface and can be equated to Dx' Dd or Dl' 1 is the length and n is the width.

The values of Dx ' Dd and Dl obtained from equations (2), (3) and (4) and, corrected for edge effects, and Yd' calculated from a rearrangement of equation (1), are given in Table 1:

Yd = (BD + 2) ([Dx/Dl] - 1)

(BD - 1) (2 + [Dx/Dl])

(6)

The volume fraction of the disperse polystyrene phase in the Kraton G-1657X model two-phase system was found by this analysis to be 0.20. Considering the nature of the assumptions made in this analysis, this value is in close agreement with the manufacturers reported value of 0.12 and demonstrates the validity of this analysis.

In the Crystic 272 polyester resin system the values of Dx' Dd and Dlobtained from equations (2), (3) and (4), using experimental values of mx and md' were not sufficiently accurate for detailed analysis. However, these values of Dx' Dd and Dl can be used for an iterative computation using equation (7) in order to obtain more self consistent values (Table 1). The values of Dd for the polyester resin, obtained using this method, are identical to the value of Dd measured for Kraton G-1657X. Thus the moisture diffusion coefficient for the dense, glassy phase in the polyester resin is similar to the polystyrene phase in the block copolymer. This is not altogether surprising since they are both polymers of styrene.

M(t) = Yd[1 - exp(-7.3 [Ddt/h2]0.75)] (7)

+ (1 -Yd) [1 - exp(-7.3 [Ddt/h2]0.75)]

This equation, derived in a previous paper/7/, enables the sorption curve for a two-phase material, normalised with respect to Mm, to be predicted from a knowledge of Dx' Dd' Dl and Yd' It is not possible to vary Yd independently of Dx' Dd and Dl and so the theoretical curves calculated from this equation will only fit the experimental data under a well defined set of conditions. It is also noted that the absorption curves of the specimens exposed to different relative humidities were identical when normalised with respect to Mm, and this demonstrates that in this system diffusion is concentration independent (Fig. 2).

The volume fraction of the dense phase calculated from equation (6) (Yd = 0.16) is lower than would be predicted from the work of

631

Bergmann and Demmler/8/. These authors suggested a structure of dense phase microgel particles of 500A radius separated by 50A of less dense material. If we assume that these microgel particles are arranged in a cubic array then Vd for the system would be 0.45. However, in such an analysis it is assumed that the dense phase is of discrete particles of uniform density, which is unlikely to be the case. The probability is that the microgel particles have a dense core, but become progressively less dense towards the periphery of the particles.

Further evidence for the two-phase structure of Crystic 272 polyester resin was supplied by scanning acoustic microscopy (SAM). Under the Leitz ELSAM acoustic microscope at a frequency of 1.8Hz, the image of the resin was dappled in appearance (Fig. 3). The lighter shade, circular regions surrounded by the darker continuous phase visible in the micrograph are believed to be regions of higher density material embedded in a less dense matrix. This is consistent with the heterogeneous structure of unsaturated polyester resins described in the literature/8-l0/ and is also consistent with the two-phase diffusion model.

CONCLUSIONS

Crystic 272 isophthalic unsaturated polyester resin has a two­phase structure comprised of particles of the dense phase embedded in the less dense phase. The moisture absorption curve for a two-phase polymeric material of this type can be described by considering the moisture absorption characteristics of the two phases. Using a Fickian approach the diffusion coefficients of the dense and less dense phases can be calculated, and from these values the volume fraction of the dense phase in the material can be calculated.

ACKNOWLEDGEMENTS

One of us (PMJ) wishes to thank the SERC for a research studentship. We thank Mr. M. Issouckis of E. Leitz (Instruments) Ltd. for help with the acoustic microscopy, and Scott Bader and Co. Ltd. for a gift of resins.

REFERENCES

1. C.H. Shen and G.S. Springer, J. Compo Mater., 10 (1976) 1

2. P. Bonniau and A.R. Bunsell, J. Compo Mater., 15 (1981) 272

3. H.G. Carter and K.G. Kibler, J. Compo Mater., 12 (1978) 118

4. A.C. Loos and G.S. Springer, J. Compo Mater., 13 (1979) 131

5. B. Ellis and M.F. Found, Composites, 14 (1983) 237

6. E. Behrens, J. Compo Mater., 2 (1968) 2

·7. P.M. Jacobs and F.R. Jones, J. Mater. Sci., in press.

632

8. K. Bergmann and K. Demmler, Koll. Ztscht. und Ztsch fur Polym., 252 (1974) 204

9. W. Funke, J. Polymer Sci., C16 (1967) 1497

10. F.R. Jones and M.J. Mulheron, Composites, 14 (1983)

Crystic 272 Kraton G-1657X

75%RH 96%RH Immersion 96%RH

Dx 8 8 8 1.89

Dd 0.16 0.16 0.16 0.16

Dl 10 10 10 2.49

Vd 0.16 0.16 0.16 0.20

:r.1i.~1.E! .... :1 .... Moisture absorption diffusion data for fully cured Crystic 272 polyester resin and Kraton G-1657X thermoplastic elastomer at 50"C. See text for definitions. All values of diffusion constants, D, are in 10-5mm2s-1

M (%)

.4r------,--------------------~--------~

.3

.2

.1

I I I I • . ~ I . ,

II

. . ... . I I

• I I I I I

o ~~----~--~--~--------------~--~~ o 200 400 800 800 1000 1200

v(tlme)/thlck ness

fi.g. 1

Experimentally determined moisture absorption curve for Kraton G-1657X.

M/M •

. 8

.6

.4

.2

OL-~~~--~--~------~--~--~~--~--~--~

o 100 200 300 400

v(tlme)/thlckness

Fig. 2

500 600

633

Comparison of the theoretical and experimental moisture absorption curves for Crystic 272 polyester resin: a) immersion (.); b) 96% RH (A); c) 75% RH (.).

Fig. 3

SAM image of Crystic 272 polyester resin.

INFLUENCE DU VIEILLISSEMENT SUR LE COMPORTEMENT AU PERLAGE DE TUBES VERRE·RESINE

ABSTRACT

I. GHORBEL, D. VALENTIN, M.C. YRIEX·, J. GRATTIER·

Ecole des Mines de Paris Centre des Materiaux - BP 87 - 91003 EVRY - France

"EDF Les Renardieres - Departement Etude des Materiaux BP1 - 77250 MORET SILOING - France

The hygrothermal aging induces physical and chemical damages into composites leading to lost of mechanical properties and change in the damage mechanism and fracture behaviour. This paper describes the mechanical behaviour under internal pressure with closed end pressure testing procedure of glass filament wound pipes. This paper reports the interaction between physico-chemical and mechanical damage mechanism and the effect of both ageing time and matrix flexibility on the mechanical properties of the tubes.

INTRODUCTION

Pour l'equipement ou le remplacement de circuits de refroidissement de centrales thermiques et nucleaires vehiculant de l'eau brute a 6 bars et 60·c, Electricite de France envisage l'utilisation de tuyauteries en materiaux composites, dont la duree de vie devra depasser trente ans. Afin de mieux etudier les mecanismes d'endommagement sous l'action du vieillissement hygrothermique, une caracterisation physico-chimique par DCS et infra-rouge a ete realisee, tandis que des essais de mise sous pression interne ont ete effectues sur deux fabrications industrielles de tubes, fabriques par enroulement filamentaire (± 55·), a l'etat de reception et apres plusieurs milliers d'heures de vieillissement dans de l'eau a 60·C.

I - MATERIAUX ETUDIES

Les deux fabrications de tubes etudies presentees dans Ie Tableau 1 se differencient essentiellement par la nature de leurs resines (fig.l) et presentent une structure similaire, a savoir :

636

- une couche interne anti-corrosion de reSlne pure renforcee d'un voile de surface et d'un mat de verre (epaisseur de 1 a 2 mm).

- une partie structurale realisee par enroulement filamentaire a ± 55' de roving de verre (epaisseur 3 a 4 mm).

Ces tubes sont dimensionnes par les fournisseurs pour une pression nominale de l'ordre de 10 bars et une duree de service de 30 ans.

II - ETUDE PHYSICO-CHIMIQUE DU VIEILLISSEMENT HYGROTHERMIQUE

2.1. Absorption des resines pures et renforcees

Une etude prealable du vieillissement hygrothermique a ete realisee sur les reSlnes pures A et B, et sur composites correspondants. A cet effet, les echantillons ont ete immerges dans de l'eau a 60'C et les cinetiques d'absorption ont ete determinees par prise de poids reguliere. Les courbes d'absorption (fig. 2) revelent Ie comportement Fickien des deux resines, alors qu'il est plutot du type Langmuir pour les composites. A masse de resine egale, on remarque que l'absorption des composites est plus importante. On peut expliquer cette absorption plus importante essentiellement par la presence d'interfaces dont la structure physico-chimique est differente. Au vu de ces resultats, on peut entre autre supposer une plus grande concentration d'eau au vOlslnage des fibres. On note au cours du vieillissement un changement d'apparence des deux materiaux s'accompagnant d'une perte de poids pour Ie composite B au bout de 3900 heures. Cette perte de poids peut correspondre a un rejet de certains composes chimiques de la matrice /1/, consequence d'une eventuelle hydrolyse.

2.2. Suivi par D.S.C. des proprietes physiques des materiaux vieillis.

Pour illustrer l'effet combine de l'eau et de la temperature sur les materiaux A et B, nous avons SU1Vl l'evolution, au cours du vieillissement, de la temperature de transition vitreuse "Tg" par D.S.C. Nous en avons deduit qu'il y a plastification des matrices A et B au cours de leurs sejours dans de l'eau a 60'c. Cette plastification se traduit par la baisse du "Tg" (fig.3) et est souvent expliquee /2/ par l'insertion des molecules d'eau dans Ie reseau macromoleculaire, conduisant a la destruction partielle de la cohesion moleculaire du reseau et a l'augmentation de la mobilite moleculaire.

Nous avons remarque, par ailleurs, a partir des courbes de D.S.C. du materiau B, l'existence d'un pic exothermique qui diminue au cours du vieillissement. Nous en avons deduit qu'en fin de fabrication, Ie composite B presente un certain nombre de chaines non pontees et que l'eau et la temperature ont active la formation de nouvelles liaisons au cours du vieillissement augmentant ainsi, son taux de reticulation.

2.3. Suivi par Infra-rouge de l'evolution chimique des composites vieillis.

L'immersion dans l'eau a 6o'c de produits a forte contenance de groupements esters peut favoriser une reaction d'hydrolyse.

Nous avons utilise afin de rendre compte de la possibilite d'une tel Ie reaction les Infra-Rouges comme outil d'analyse.

637

Pour mettre en evidence toute transformation chimique eventuelle, nous avons etudie les intensites relatives des principales bandes caracteristiques des deux materiaux susceptibles d'evoluer au cours du vieillissement, en utilisant la loi de Beer et en prenant comme reference la bande situee a 1500 cm- l .L'indexation des pics a ete realisee a l'aide des travaux de Bellamy /3/, et d'autres etudes concernant les polymeres /4/, /5/. Les echantillons analyses par infra-rouge se presentent sous forme de pastilles preparees a partir d'un compactage sous vide d'un melange de poudre de KBr et du materiau a analyser.

Les spectres realises changement d'intensite et (Fig.4). Cette evolution vieillissement est long. reaction ont eu lieu :

sur Ie materiau A vieilli, revelent un de forme de certaines bandes de vibration est d'autant plus frappante que Ie temps de Ainsi nous remarquons que deux types de

- formation d'un alcool primaire par reaction d'un residu de styrene avec de l'eau induisant la disparition des pics a 1646, 1410, 985 et 735 cm- l , l'elargissement de la bande de vitration de -CH2 - a

3033 et 1250 cm- l et l'augmentation de l'intensite du pic a 1050 cm- l . (v (C-O) de l'alcool primaire).

- hydrolyse de l'ester detectee par la formation d'un acide dont les bandes caracteristiques se situent a 3080 cm- l et 832 cm-l,ainsi que par la presence du groupement phenol a l'origine de l'augmentation des bandes se situant a 770 et 632 cm- l .

A partir des spectres obtenus (Fig.5), nous remarquons la coexistence de differents types de transformations chimiques subies par Ie materiau B au cours de son vieillissement a savoir :

- reactions de reticulation et d'hydrolyse pendant les premieres 4000 heures. L'evolution est caracterisee par la diminution de pics de vibrations des liaisons caracteristiques du styrene (3059, et 959 cm- l ) et du vinyl en fin de chaine (3034, 830 et 554 cm- l ). Ces evolutions sont d'autant plus importantes que Ie temps de vieillissement est long. L'hydrolyse est mise en evidence par la diminution de certaines bandes de vibrations du groupement ester (1130 cm- l ), par la formation d'un acide responsable du dedoublement du pic se situant a 1450 cm-let par la formation d'un alcool primaire.

- dans un deuxieme temps, nous remarquons qu'au dela d'une certaine duree de vieillissement (> 4000 h), il y a formation d'un groupement epoxy par desydratation, a l'origine de l'augmentation de certaines bandes se situant a 1250, 1130 et 520 cm- l .

III EVOLUTION DES CARACTERISTIQUES MECANIQUES DES COMPOSITES TESTES A DIFFERENTS TEMPS DE VIEILLISSEMENT.

Les tubes vieillis ont ete testes en pression interne avec effet de fond et les comportements observes ont ete compares avec ceux des tubes non vieillis. Nous avons sui vi Ie comportement mecanique en temps reel par un extensometre a jauges, capable de mesurer les deformations longitudinales et diametrales. L'emission acoustique (nombre de coups cumules) correlee aux observations micrographiques a permis Ie suivi de l'endommagement des materiaux testes: Ie mode de perissement considere etant Ie perlage.

638

3.1. Description des mecanismes d'endommagement

Le perlage des deux tubes consideres s'operent en deux temps : - Fissurat~on de la partie mecanique consistant en une fissuration

intralaminaire consecutive a une decohesion interfaciale, puis fissuration interlaminaire. Les seuils d'apparition de ces mecanismes sont differents de par les allongements a rupture des matrices respectives.

- Rupture du liner provoquant la fuite du tube et eventuellement eclatement de celui-ci.

A l'etat de reception la degradation de la partie mecanique du tube A s'opere par fissuration homogene conduisant lors de la rupture du liner a un pleurage bien reparti sur tout Ie tube, et caracterise par l'apparition de gouttes d'huile a sa surface. Au cours du vieillissement une modification du mode d'endommagement s'opere : la densite des fissures diminue tandis que leur longueur augmente et des delaminages apparaissent. Contrairement aux tubes A, a l'etat de reception, les tubes B ont un perlage tres localise, caracterise par un jet continu, et pour des press ions plus elevees. Cette difference s'explique par la difference de ductilite des matrices. Ces resultats sont similaires a ceux obtenus par Legg et Hull 161 et Carswell 17/,par contre, si les non-linearites observees sur les courbes pression-deformation (Fig.6) sont correlees avec l'apparition des prem~ere~ fissures, il faut les attribuer aussi au comportement non-lineaire de la matrice pour Ie composite B.

Au cours du vieillissement, on assiste a une modification du mode de perissement de la structure B, puisque, si a partir de 1000 heures en immersion, la rupture par eclatement est legerement precedee d'un pleurage, au-dela, on observe toujours un perissement par rupture de fibres qui ne peut ~tre seulement explique par les modifications physico-chimiques de la matrice, mais eventuellement par la degradation des fibres de verre 18/.

3.2. Influence du vieillissement sur Ie comportement mecanique des tubes.

Pour les deux tubes, on cons tate (Fig.7), une baisse de la pression de perlage en fonction de la duree de vieillissement. Cette evolution se stabilise au bout de 4000 heures pour Ie tube B, mais semble se poursuivre pour Ie tube A. On remarque toutefois que les niveaux de pression obtenus res tent tres superieurs a la pression nominale de service des tubes. La tolerance a l'endommagement des tubes peut etre mesuree par Ie facteur k = Pp/Pe, rapport de la pression de perlage sur la pression seuil de fissuration, determinee par apparition d'une non-linearite sur les courbes P = f (€). On cons tate (fig. 8), un comportement oppose pour les deux tubes, k augmente nettement pour Ie materiau A alors qu'il semble diminuer pour Ie materiau B. On peut traduire cette evolution par une baisse du seuil de fissuration des tubes A plus rap ide que la baisse de pression de perlage, contrairement au tube B, pour lequel Ie seuil de fissuration se rapproche de la pression de perlage expliquant la rupture catastrophique des tubes. On remarque, la encore, une stabilisation de cette evolution vers 4000 heures. L'evolution du nombre de coups

639

cumules, (fig. 9), montre une diminution de l'endommagement accumule necessaire pour arriver au perlage, et qui se stabilise plus rapidement pour Ie composite A. Ceci est a rapprocher des observations micrographiques montrant une fissuration moindre de la partie mecanique.

IV DISCUSSION

Le vieillissement hygrothermique a donc entraine un changement physico-chimique des composites se manifestant par une plastification des matrices contrebalancee dans Ie cas du composite B par l'achevement de la reticulation qui peut conduire a une rigidification supplementaire. Par ailleurs d'autres reactions chimiques peuvent avoir lieu, telles que l'hydrolyse des deux matrices et la deshydratation de la matrice B, par formation de groupements epoxy au profit des fonctions hydroxyles des alcools primaires et secondaires. Du point de vue mecanique, la diffusion d'eau entraine la creation de contraintes internes et une degradation de l'interface par infiltration de l'eau au sein des vides aces endroits, les deux mecanismes conduisant a la decohesion de 1 'interface, expliquant pourquoi la quantite d'endommagement cree Par la mise sous pression necessaire au perlage est moindre. La degradation des caracteristiques residuelles des composites vieillis s'opere en deux temps:

- une degradation rapide des proprietes pour des durees inferieures a 1000 heures pour Ie composite A et a 3000 heures pour Ie composite B.

- une evolution lente pour Ie composite A, une quasi-stabilisation pour Ie composite B.

V CONCLUSION

Cette etude a montre l'interaction complexe entre la degradation mecanique et physico-chimique des composites vieillis.

Nous avons pu cons tater que Ie vieillissement dans de l'eau a 60'c conduit a une modification des parametres rheologiques des matrices entrainant une modification du mode d'endommagement et de rupture.

Toutefois les degradations des interfaces et de la matrice dues aux vieillissement creent une surcharge sur les fibres qui, s'ajoutant a leur degradation, modifient Ie mode de rupture de ces tubes apres quelques milliers d'heures dans de l'eau a 60'c; on assiste a des ruptures par eclatement au lieu du perlage cons tate a l'etat de reception. L'etude se poursuit par une analyse du comportement en fluage de ces deux fabrications de tubes.

REFERENCES

1 - Jacquenet R., Lagrange A., Vieillissement de stratifies polyester/verre E et evolution de leurs caracteristiques mecaniques en milieu marin". Colloque COMPOSITES-88, 20-22 juin 1988, Nice, COMPOSITES STRUCTURES, pp. 260-277.

2 - Verdu J., Action de l'eau. Techniques de l'Ingenieur, A 3165-1 A3165.

3 - Bellamy L.J., The infrared spectra of complex molecules. Volume One. Third Edit. Chapman and Hall.

640

4 - Kemp W., Organic spectroscopy. pub1ie par The Nacmi11an Press LTD.

5 - Handbook of chemistry and physics. 62 ND Edition, (1981-1982). 6 - Legg H.J., Hull D., Effects of resin flexibility on the

properties of filament wound tubes". Composites, October 1982, pp.369-376.

7 - Carswell W.S., The behaviour of glass filament wound pipes under internal pressures. Proc. Ind. Int. Conf. on Composite Materials (1978), pp. 472-483.

8 - Ishai 0., Environmental effects on deformation, strength and degradation of unidirectionna1 CFRP. Po1ym. Eng., Vol. 15, n" 7, (July 1975) .

Fibre de Nombre porosite Tg Resine € % verre de Vf % ( "C)

matrice utilisee couches % tube A polyester 0,7 E 8 49 2 132 tube B viny1ester 5 ECR 10 54 1 110

Tableau I : Caracteristique des tubes a l'etat de reception.

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l J ME. kctur . ·· AGEING TI ME <hrs)

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. : - CH=~­

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641

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MOISTURE ABSORPTION INFLUENCE ON THE MECHANICAL PROPERTIES OF CARBONIEPOXY COMPOSITES

ABSTRACT

I. MONDRAGON, J. IMAZ', A. RUBIO', A. VALEA

EU/TI- P.Pio XII sin - SAN SEBASTIAN - Spain */NASMET - BO /gara sin - SAN SEBASTIAN - Spain

The different properties of the high performance composite materials may suffer certain modifications as.a result of

environmental effects. This report deals with the influence of the hygrothermal conditions over the static mechanical properties of a laminated made of an epoxy resin reinforced with carbon fibre. Previous to the mechanical characterization of the material, the kinetic of the moisture absorption under different conditions has been studied. This lead to a relationship between the modifications suffered by the laminate and the variation of the mechanical properties.

INTRODUCTION

The applications of the composite materials as structural materials are increasingly more relevant. The development of such applications requires a careful knowledgment of both the environment where the material may be found and the influence of the environment over its mechanical properties. At this level main factors are environment and temperature, together with a combination of both.

For above mentioned reasons, it becomes necessary to simulate the influence of both factors when intending to reach a real approach of the degradations study produced by environment. The researchs mentioned in different papers show that when composite laminates are exposed to certain hygrothermic pressure, moisture is absorbed over the resin and it diffuses through the resin until an equilibrium is theoretically reached after a long time of exposure. This absorbed moisture may reach the interphase between the fibre and the matrix, which produces plastification of the matrix affecting the properties of the material./l/,/2/,/3/.

644

Main objective of this work is to study the kinetics of the moisture absorption over a certain laminate besides its variations with several environmental external conditions. At the same time, changes observed in connection with the modification of the mechanical properties suffered by the material are also investigated.

MATERIALS AND EXPERIMENTAL PROCEDURES

The material employed for performing this task was an epoxy resin laminate of DGEBA, reinforced with carbon fibre. The T-300 carbon fibre laminates displayed (0,90)6. The weight ratio of fibre is 68%. The laminate used was manufactures from plates preimpregnated by means of the hot compression method. The nominal thickness of the obtained laminate was about 2 mm from which specimens were further mechanized. Dimensions of specimens employed were 50x25x2 mm. Before introducing the specimens in the wet environmental treatment(liquid or moistured atmosphere), these were dried in a dissecator at 20~C and with a silica-gel until no weight variation was observed in the specimens.

Specimens were placed in the proper hygrothermal environments measuring their weight by means of periodical weighing in a Metler Analytical Balance. Environments employed in the hygrothermic treatments are displayed in Table 1. Higher temperatures were not employed in order to avoid interaction with either the curing effects or the degradation of properties to be measured.

~ 20 35 50 65 lance 6e% x x x x

75% x x x x

100% x x x x

Immerg. x x x x

Table 1. Treatment conditions

Curves representing the quantity of absorbed water on the specimen versus treatment time were obtained by periodical weighing. The weight variation on the material may be described as:/4/

Mt MoO

Specimens were characterized mechanically in both the initial product and after the hygrothermal treatment. Mechanical characteri zation was achieved by flexural test according to ASTM D-790 and

645

stress-strain curves were obtained in a Instron 4026 device.

RESULTS AND DISCUSSION

Results displayed in Fig 1 show the weight variation (i.e., moisture absorption over specimens as a function of the time exposure). Such result suggest that the moisture absorption process varies as follows:

a) Depending on the environment,i.e., either by water inmersion or by atmosphere moisture ratio.

b) Depending on the temperature The moisture absorption curves referred to a certain atmosphere exposure display a fickian behaviour with a straight initiation followed by a stabilization once the equlibrium between the environmental humidity and the water absorbed by the specimen is reached. However, such equlibrium during time exposure was not totally achieved by specimens inmerged in water.

By studying the different curves of fig 1, it may also be observed the strong dependence on temperature of the kinetical absorption process, while the water quantity absorbed to equilibrium seems a parameter which depends more on the water concentration of the environment than on the temperature itself.

The diffusion coefficient may be obtained from the initial section of the curve by means of using those equations from the fickian behaviour.

The influence of the absorbed moisture over the bending strength has been displayed in Fig. 2. It may be observed from such figure how mechanical properties have diminished as a result of this absorption.

CONCLUSION

Assuming the different results obtained the following conclusions may be observed: - Estimates of the weight variation of the specimen until the

absorbed moisture saturation has been reached may be obtained in almost all cases by the Fick law.

- In the temperature and moisture ranges quantity of moisture to saturation depends on the moisture concentration of the environment while the absorption kinetic depends on temperature.

- The absorbed moisture leads to an appreciable reduction of the mechanical properties of the material measured trough the bending strength.

646

1. 25

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647

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% Absorbed moisture

Fig. 2 Residual flexural strength vs. absorbed moisture percentage

REFERENCES

1. Bonniau P. & Bunsell A.R. "A comparative study of water absorption theories applied to glass epoxy composites". Journal of Composite Materials. Vol. 15, (1981), 272.

2. Loos A.C. & Springer G.S. "Moisture absorption of graphite-epoxy composites inmersed in liquids and in humid air". Environmental Effects on Composite Materials. Chap.4. Ed: G.S. Springer. Technomic P.C. Inc. (1981)

3. Boll D.J., Bascom W.O. & Motiee B. "Moisture absorption by structural epoxy-fiber composites". Composites Science and Technology, 24, (1985), 253-273

4. Crank J. "The mathematics of Diffusion" Claredon Press, Oxford (1956) p 42

FRACTURE

Chairman: Dr G. Oi ORUSCO Montedison

THE INFLUENCE OF FIBER BUNDLE REINFORCEMENT ON THE FRACTURE MECHANICAL

BEHA VIOR QF POL YCARBONATE AND EPOXY

J. KREY, K. FRIEDRICH, K.H. SCHWALBE'

Technische Universitat Hamburg-Harburg Harburger Schlossstrasse 20 - 2100 HAMBURG 90 - West Germany

'GKSS-Research Center Max-Planck-Strasse - 2054 GEESTHACHT - West Germany

ABSTRACT

This study deals with the fracture mechanical behavior of composite materials consisting of different kinds of single fiber bundles embedded in a polymeric matrix. The fracture toughness and the failure mechanisms of these model composites are discussed in terms of fiber volume fraction and fiber/matrix interfacial bond quality. The reinforcement effect of the embedded fibers is described by a reinforcement effectiveness parameter "RMV". The normalized fracture toughness of the EP-composite systems increases strongly with increasing RMV, whereas an opposite trend was observed for the PC-composites. An improvement in toughness of the PC-composites could be achieved by impregnating the fiber bundles with a PC-solution before embedding.

INTRODUCTION

The addition of fibers to a polymer matrix results in a composite material, which may possess a very anisotropic property profile. Especially, mechanical properties, fatigue crack propagation and fracture toughness of this kind of material strongly depend on (i) the fiber orientation, (ii) the mechanical properties of the components, (iii) the fiber volume fraction, and (iv) the bond quality of the fiber/matrix interface 11/.

Model composite systems conSisting of single fiber bundles embedded in a polymeriC matrix were chosen to perform fundamental investigations on the interactions between a propagating crack, a polymeric matrix material and a single fiber bundle.

652

I - MATERIALS, MANUFACTURING PROCEDURES AND TEST CONDITIONS

1.1. Materials

To focus on extreme conditions, two matrices of different toughness, i.e. ductile thermoplastic polycarbonate (PC) and a brittle thermosetting epoxy resin (EP) were used. As reinforcement three different types of fiber materials -glass-fibers (GF), carbon-fibers (CF) and aramid-fibers (AF) - served. The influence of the fiber volume fraction on the reinforcement effect was simulated by different amounts of filaments in the bundles (roving strength FRS = load carrying capacity of the whole fiber bundle). FRS varied from 370, 740 to 1400 N.

1.2. Manufacturing procedure of the model composite systems

Two different procedures were used to manufacture the model composite plaques. The EP and EP-composites were produced by casting and the PC and PC-composites by hot pressing. PC-composites have been investigated with and without a preimpregnation of the fiber bundles in a 25%­PC-solution before embedding.

1.3. Test procedures

1.3.1. Fiber/matrix bond quality

For the characterization of the bond quality, especially the chemical adhesion between the fibers and the matrices, single filaments were embedded and the critical fiber lengths Ie were determined. Additional modified ·pull out" tests on the model composite systems showed the influences of the amount of filaments in a bundle, the matrix shrinkage and the viscosity of the matrix on the fiber/matrix bond quality. Four different test procedures were performed - three point bend-, fiber bundle tensile-, shear­and impact-tests - in order to characterize the fraction of the fiber bundle with a good adhesion to the matrix material (t1/tO) and the critical fiber length of the whole bundle.

1.3.2. Fracture toughness

Fracture toughness testing was performed at a constant strain rate in order to determine the stress intensity factor Kc' All materials were tested at room temperature and a cross head speed of 0.5 mm/min. Three point bend (TPB) specimens with a specimen width of W = 12.7 mm and CT-specimens with W = 50.8 mm were tested. In order to characterize the influence of the fiber bundle position on the fracture toughness and the fracture mechanisms two different types of CT-specimens were studied. KcA represents the K-values of specimens of type A (fiber bundle directly in front of the crack tip), KcB those of type B (fibers positioned 8 mm behind the initial crack tip ao) and Kc those of the TPB-specimens.

653

II - RESULTS AND DISCUSSION

2.1. Fiber/matrix bond quality

The results of the single fiber tensile tests are presented in Fig.1. The EP-CF-combination achieves the highest interfacial bond strength 'ti' whereas the one of the same kind of fiber embedded in the PC-matrix is only about 40% of this value.

Due to the high viscosity of the PC-matrix during the manufacturing process the embedded fiber bundles are poorly impregnated. The bond quality decreases with increasing amount of filaments in a bundle.. An improvement of the fiber bundle/matrix interfacial bond quality can be achieved by impregnating the bundles before embedding. In these systems, however, no significant influence of the roving strength (fiber volume fraction) can be detected.

The fiber bundle/matrix interfacial bond strength of all EP-composite systems decreases with increasing roving strength, but all values are in a higher range compared to the PC-systems. Due to the low viscosity of the EP­matrix a good impregnation of all fibers in a bundle can be achieved. Only the relatively high shrinkage of the EP during the curing process results in a reduction of the bond quality. Especially in the EP-CF-system this effect of fiber/matrix debonding due to the matrix shrinkage can be observed.

2.2. Fracture toughness

The determination of the fracture toughness of the model composites is based on the concepts of the linear elastic fracture mechanics (LEFM) /2/. Its application seems to be doubtful, especially in those cases, in which the final fracture of the samples occurs parallel to the embedded fibers (CT-specimens of type A and B of the systems EP-CF, EP-GF and EP-AF with the highest roving strength). However, the stress intensity data can be used to characterize the reinforcement effect of a single fiber bundle in a composite and to compare this effect to other composites, produced in a similar way.

The average fracture toughness (average of the data from the CT- and TPB-specimens) of the ductile PC-matrix is in the range of Kern = 3.7 MPa.Jm and that of the brittle EP-matrix at about Kern = 2.4 MPa.Jm.

Fig.2 represents the results of the PC-composites, where the average fracture toughness data (Kee) of KeA, KeB and Ke are shown as a function of roving strength. The reinforcement effect in the PC-systems without a fiber bundle preimpregnation is not very pronounced. The Kee-values of the impregnated PC-systems are on a higher lever, but in both composite systems no significant influence of the fiber volume fraction and the fiber position can be detected. It is assumed, that this is mainly due to the high amount of plastiC deformation of the matrix in front of the crack tip, which causes fracture of the brittle fibers before the deformed polymer finally separates. The result is the

654

same failure sequence in all PC-composites - first fiber/matrix debonding, then fiber cracking and pull out /3/.

In Fig.3 the Kcc-values of the EP-composite are presented as a function of FRS' The fiber bundles cause a clear improvement in fracture toughness with increasing fiber volume fraction. In general, two different failure mechanisms can be observed in the CT-specimens in addition to matrix fracture, either fiber cracking or matrix cracking along the embedded fiber bundle. But the fracture mechanism is independent of the kind of reinforcement and its position.

In the following an attempt is made to modify the "reinforcing effectiveness parameter" R /4/, which is valid for short fiber reinforced composites, and to discuss the normalized fracture toughness data (K' = (Kcc -Kcc,R=O)/Kcm) of the model composites as a function of RMV' The reinforcement effect, expressed by RMV, is a function of the fiber volume fraction, which is characterized by the roving strength and the fraction of the embedded fiber bundle with a good adhesion to the matrix (FRS' t1/tO)' and the fiber/matrix interfacial bond quality, which is represented by the ratio of the composite shear strength to the matrix shear strength (tdtm)' determined in the shear tests:

(1 )

In Fig.4 the normalized fracture toughness data of all model composites are presented in function of the above defined RMV ' With respect to these parameters a clear increasing tendency of the fracture toughness of the EP­composites with increasing RMV can be observed, which is due to good impregnation of the fiber bundles and the relatively high interfacial bond quality.

As a result of the weak interfacial bond quality and the high amount of matrix deformation in front of the crack tip the slope of the K' - RMV - slope of the PC-systems without fiber preimpregnation is negative. An improvement can be achieved by embedding fibers of high stiffness (e.g. CF) or by preimpregnation of the whole fiber bundle before embedding, which results in a better bonding of the fiber bundle to the matrix.

III - CONCLUSIONS

In conclusion, the following findings of this study should be emphasized: 1. An improvement of the bond quality between a fiber bundle and a matrix of high viscosity can be achieved by a preimpregnation of the bundles before embedding. 2. The K concept was used to characterize the reinforcement effect of single fiber bundles on the fracture toughness of polymeric materials. 3. A correlation between the microstructural parameters - fiber volume fraction and fiber/matrix bonding - and the fracture toughness can be expressed by a modified reinforcement effectiveness parameter RMV'

655

4. The fracture toughness of the model composites can be improved by using (i) reinforcing material of high stiffness, (ii) a brittle matrix and (iii) preimpregnated fiber bundles in matrix materials of high viscosity.

REFERENCES

1. M.J. Folkes, Short Fibre Reinforced Thermoplastics, Research Studies Press, Chichester (1982)

2. K.-H. Schwalbe, Bruchmechanik metallischer Werkstoffe, Hanser, MOnchen Wien (1980)

3. J. Krey, K. Friedrich, K.-H. Schwalbe, J. Mater. Sci. Letters 6 (1987) 851-855 4. K. Friedrich, Compo Sci. Techn. 22 (1985) 43-74

~y-----------------~-----------,

EP·COMPOSITES PC·COMPOSITES

Fig.1 - Interfacial bond strength of the single filaments in EP and PC 10.0

r;] PC-CF A PC-GF

7.5 : PC-Cf ..... ,

PC-GF ......

f &? -:!.

5.0 u ..A. u

>< ~ r---.. -~

2.5

0,0 o 500 1000 1500 2000

Roving Suenglh F RS [N]

Fig.2 - Fracture toughness of the PC-model-composites as a function of the roving strength

656

f 'E ~

7,5

• EP-CF

o EP-GF

~ 5,0 +-----+-~:..-..,.e.:e,.......::::::..--+---_I u u ~

2,5 +-----+-----1----+----1

0,0 -I--........ ---+.----.--"--"T"'""--I----.----I o 500 1000 1500 2000

Roving Strength F AS [N) -

Fig.3 - Fracture toughness of the EP-model-composites as a function of the roving strength

4

• EP-CF o PC-CF

1 3 0 EP-GF l:J. PC-GF

" EP-AF • PC-CF,impr.

2 A

0 , a: 8 Ii ~~ 8 0 ~

-1

-2 0 200 400 600 800 1000 1200

RMV ~

Fig.4 - Normalized fracture toughness of the composite systems as a funktion of RMY

ABSTRACT

TRIGGER MECHANISMS IN ENERGY ABSORBING GLASSCLOTH·EPOXY TUBES

I. SIGALAS, M. KUMOSA", D. HULL"

Council for Scientific and Industrial Research Division of Materials Science and Technology, CSIR

PO Box 395 - 0001 PRETORIA - South Africa 'University of Cambridge. - Department of Materials

SCience & Metallurgy - CB2 30Z CAMBRIDGE - England

The triggering mechanisms in energy absorbing glass cloth-epoxy tubes were investigated using microscopic examination supported by finite element analysis.

It is found that the initial stages of the crushing process are dominated by the continuous formation of wedges that generate lateral cracks that cut off small rings of material. This process is at a later stage modified to a front-wedge-front geometry.

1 - INTRODUCTION

The interest in the use of composite materials for energy absorp­tion purposes stems from crash worthiness conSiderations in the car and helicopter manufacturing industry [1-7].

Maximum energy absorption is achieved when the composite mate­rial, which is glass or carbon fibre reinforced reSin, in some form of tubular structure, is made to crash in a stable mode by propagating a crash zone through its length. In order to achieve this, various trigger mechanisms have been used, the aim being to generate various stress concentrations at the end of the tube thus promoting stable crush rather than brittle fracture of the whole structure [8-13].

It has been found in this laboratory that depending on the trig­gering mechanism used, different loads are required to initiate the stable crash process. This variation can be of use in the car or hel-

658

icopter designer, as it allows him to tailor the performance of his energy absorbi ng structures. Despite thei r importance, triggeri ng mechanisms are very poorly understood. Up to now very little work has been done in studying the micromechanics of the triggering process.

In this work we report our study of trigger mechanisms of glass­cloth-epoxy tubes at the microscopic and macroscopic level, with par­ticular emphasis on chamfer triggers.

2 - EXPERIMENTAL RESULTS AND DISCUSSION

Figure l(a) shows the tube geometry used. In order to explain the crushi ng processes as observed in the mi crographs obtained, we used finite element modelling. The effort required that the elastic constants and mechanical strengths of the material out of which the tubes had been made be measured. Consequently, the main body of work in this investigation was preceded by a comprehensive progranme of material characterization and mechanical properties measurement.

The tubes were obtained from Tufnol Ltd., UK and they were of the type RLG1. They were SO nm in internal diameter and 2.S nm in wall thickness. Flat sheets, 3 and 2S nm thick of the same material were used to measure in plane and through plane properties, respectively.

Tensile tests in the 1,2 and 1,3 planes were used to generate E1, v12' v13 and 01*. The losipescu test was used to generate values for G12 , Gu , t12* and t13*. A transverse compressive test was used to generate values for E3 and v31. Table 1 shows the values obtained by using the above tests. The fibre content was 48% by volume, while the void content amounted to S%.

The tube testing progranme involved the crushing of tubes with chamfer angles of 10°, 20°, 30°, 40°, SOo, 60°, 70°, 80°, 8So and 90°. In each case the first tube was crushed to the point where the crush zone had fully consumed the chamfer region and an average crush load had been established. The distance covered by the mechanical testing machi ne ram ·from the poi nt of fi rst contact with the tube's end, to the point where the average load had been established was noted. The next tube was then crushed until the ram displacement after first con­tact had attained 0.1 of that distance. The next tube was crushed until the ram displacement attained 0.2 of that same distance and so on until a total of 10 tubes had been used. The tube's length, from the beginning of the chamfer, to the bottom of the tube was SO nm for all tubes tested.

It was found that the crushing load attained an initial value, which we shall call triggering load, and which was different from the average load established after sufficient ram motion (typically 12 nm). The difference between trigger load and average load was found to be a pronounced function of the chamfer angle, it being high and positive above 70° while being negative at around 45°.

659

The tubes obtained in the manner described above were cast in a clear resin under load. They were then dissected, the exposed surface polished and the whole crushed region photographed to a total magnifi­cation of x30. The following features became apparent when examining the micrographs:

For chamfer angles of 10° and 20°, fibre bending was apparent, although brittle fracture of the material was taking place. Large sca 1 e bendi ng does not occur beyond tube wa 11 thi cknesses of 0.6 mm. From that point on the material crushes by means of lat­eral cracks moving either in or out of the tube, always downwards, and at 30° to the tube's axis. For chamfer angles larger than or equal to 30°, crushing was ini­tiated by slight bending of the material towards the inside of the tube. This gave rise to the generation of circumferential cracks, pointing towards the inside of the tube and at 30° to the tube's axis, as shown in Fig. l(b). The above process was repeated for several times, until, in the case of chamfer angle of 30°, the cross-section became as shown in Fig.l(c). In the case of chamfer angles higher than 30°, the initial inwards breakage would be followed by a circumferential crack, starting at the top inside end of the tube and moving out­wards at 30° to the tube's axis, as shown in Figs 2(a) to 2(c). It is the generation of this crack that requires the high trigger­ing loads occurring at high chamfer angles. It can be seen that the tube cross-section shown in Fig. 2(c) is the same as that shown in Fig. l(b). As a result of this sequence of events the tube cross-section ends up looking like Fig. l(c) irrespective of chamfer angle. At that point the material on the apex shown was compressed into the underlying zone, which was forced to deform laterally in order to make room for the intruding matter. This resulted in the for­mation of a wedge, as shown in Fig. l(d). The stress concentration at the tip of the wedge then generates either lateral cracks, or a central Mode 1 crack, depending on the relative material strength and stress distribution in the two directions. This situation is portrayed in Fig. l(e). The mode of lateral cracking shown here creates a self-correcting mecha­nism, whereby the tip of the apex is always shifted towards the centre line of the tube wall by virtue of the fact that the lat­eral crack always pOints towards the nearest tube surface. This mode of crushing was seen to be favoured in the tubes tested in this investigation.

This process carries on until the tip of a wedge finds itself at close proximity to a void. The void acts as a Mode 1 intrawall crack initiation pOint. Because of the low strain energy release rate asso­ciated with such a crack the critical crack length is very low compar­ed with that of a lateral crack. As a result the crack in question will easily become of critical length and will then carryon propagat­i ng itse If. Th is event wi 11 change the crush i ng process from one of wedge-lateral crack to the more often observed [14] wedge-front-wedge one.

660

Finite element modelling has been used to explain the following:

The initial inwards cracking, as shown in Fig. 2(b). The JO° external circumferential crack, as shown in Fig. 2(c). The compressive stresses that generate the wedge, as shown in Fig. 2(d). The stresses that generate the cracks at the tip of the wedge. A slight preference was found for the cracks emanating from a cen­trally positioned wedge to advance towards the outer tube wall. This would shift the apex of the crash zone towards the inside of the tube to a limi ted extend, in agreement with A. Fa i rfull' s findings [15].

:--.- chamfer angle

(a) (b)

o (c) (d)

(e)

Fig. 1 - Schematic presentation of tube crushing stages for chamfer angles equal to or higher than JO°

661

(0) (b) (c)

Fig. 2 - Schematic representation of the generation of the large lat­eral crack for angles higher than 30°

Table 1 - Mechanical properties of Tufnol glasscloth-epoxy material

Ell E22 27.5 GPa E33 12.03 GPa v12 v21 0.145 v13 0.40 v31 0.196 G12 = G21 5.15 GPa G13 5 GPa 01* = 02* 427 MPa 112*= 121* = 83 MPa 113* 64.7 MPa

REFERENCES

1. C.L. Magee and P Thornton, Trans. SAE, 87 (1987) 2041.

2. P.H. Thornton, H.F. Mahmood and C.L. Magee, in "Structural crush­worthiness", (Jones and Wierzbicki, ed.) (1983) 96-117, Butterworth, Boston.

3. P. Beadmore and C.F. Johnson, Compo Sci. Tech., 27 (1986) 251.

662

4. H. Vogt, P. Beadmore and D. Hull, in "Kunstoffe al s Problemloser in Automobilbau" (1987) VDI Conference, Mannheim.

5. J.D. Cronkhite, T.J. Hass, V.l. Berry and R. Winter, Investigation of the crash impact characteristics of advanced airframe structure, USARl-TR-79-11, Sept. 1979.

6. J.D. Cronkhite and V.l. Berry, USA AVRADCOM-TR-82 D-14, Feb. 1983.

7. G.l. Farley, "Energy absorption of composite materials and structure", 43rd American Helicopter Soc. Annual Forum, May 1987.

8. D. Hull, "Axial crushing of fibre reinforced composite tubes", 'Structural Crashworthiness', (Jones and Wierzbicki, eds) (1983) 118-135, Butterworths, london.

9. G.l. Farley, J. Compo Mater., 17 (1983) 267.

10. P.H. Thronton, J. Compo Mater., 13 (1979) 247.

11. P.H. Thornton and P.J. Edwards, J. Compo Mater., 16 (1982) 521.

12. P.H. Thornton and P.J. Edwards, in ICCM-V, (Harrigan, Strife and Dhingra, eds) (1985) 1183, Metallurgical SOCiety, Pensylvania.

13. D Hull, "Energy absorption of composite material s under crash conditions", Progress in Science and Engineering of Composites Vol. 1 (Hayashi, Kawata and Umekawa, eds) (1982) 861-870, ICCM-V, Tokyo.

14. A.H. Fairfull and D. Hull, "Effects of specimen dimensions on the specific energy absorption of fibre composite tubes", (F.l. Matthews et ale eds) (1987) 3.36-3.45, Sixth Intern. Conf. on Composite Materials ICCM-VI, Elsevier.

15. A.H. Fairfull, "Scaling efects in the energy absorption of axially crushed composite tubes", (1986) PhD Thesis, University of liverpool.

HYBRIDES HYBRIDS

Chairman: Pr I. CRIVELLI-VISCONTI Universita di Napoli

ABSTRACT

THE EFFECT OF AGGLOMERATION AND THE RESIDUAL STRESS STATE

ON THE PERFORMANCE OF GRADED PARTICULATE HYBRID GLASS FIBRE COMPOSITES

F. JONES, S. AHMED

University of Sheffield - School of Materials Northumberland Road - S10 2TZ SHEFFIELD - England

The micromechanics of model DO/Sand/Do glass fibre hybrid laminates has been studied. The tensile moduli are closer to the upper rather than lower bound laws of mixtures predictions. The enhanced moduli associated with particle agglomeration in resin castings is maintained in the composite. The microdamage is shown to be analogous to transverse cracking in 0°/90% 0 composite by application of the shear-lag theory.

1 - INTRODUCTION

Particulate composites are receiving considerable attention for various engineering applications since they provide improved mechanical performance economically. Fillers such as sand particles are similarly incorporated into the body of industrial glass fibre composite pipes and linings while maintaining corrosion resistance. In order to study the micromechanics of these industrial composites a DO/Sand/0° laminate configuration has been developed which is analogous to the filament wound type/I/. In this paper we have applied the shear-lag model to the effect of thickness on the micromechanics of inner ply failure.

2 - EXPERIMENTAL

The resin, Crystic 272 (Scott-Bader & Co Ltd) is an isophthalic unsaturated polyester resin which was cured using 2phr of 50% methyl ethyl ketone peroxide solution (catalyst M) and. 0.25 to 0.5 phr of cobalt naphthenate solution (accelerator A). At this concentration of

666

accelerator the resin has a relatively long gel time which enables the manufacture of void free laminates. The glass fibre used in these laminates was Silenka 600 tex E-glass fibre with a polyester compatible silane finish. The sand filler was supplied by Stanton PIc in three grades; coarse, medium, and fine. The sieve analysis showed that >94% of particles were in the range of 600 to 1180 pm, >95% from 600 to 300 pm, and 300 to 150 pm for coarse, medium, and fine grade respectively. The details for the fabrication of model hybrid glass fibre OO/Sand/Oo laminates shown in Fig 1 and cast plates of filled resin can be found elsewhere /2,3/. The radius of curvature of assymetric OO/Sand beams was used to determine the thermal strain which is built into the sand lamina of the OO/Sand/Oo composite after post curing /4/. Aluminium alloy(HS15) end tabs 30x21x1.5 mm, were bonded to all tensile test specimens with cold setting Araldite adhesive (Ciba-Geigy PIc) prior to post-curing at 1200 C for 4h in an air circulating oven. The specimens were allowed to cool down inside the oven for ~10 hrs prior to testing. 90 mm long electrical resistance strain gauges were bonded to all specimens using cyanoacrylate adhesive, in order to monitor the strain along the length of the specimen during tensile testing. The stress was calculated from the applied load and the original average cross sectional area. Tensile tests were performed on a Mayes SM200 which was calibrated to BS1610. The cross head speed was 0.5 mm/min. The modulus was obtained from the linear portion of the stress-strain curve.

3 - RESULTS AND DISCUSSION

Fig 2, shows a typical stress-strain curve for OO/Sand/Oo laminate. The curve has an initially linear portion with a rapid change in slope after a strain of 0.06%. In order to investigate these events experiments with low Vs were carried out /5/. It was observed that particle matrix debonding and tranverse matrix cracking between particles occurred. The degree of non-linearity was found to increase with increasing inner ply thickness. An irreversible whitening effect was also observed which was coincident with the initiation of damage in the inner sand lamina. The strain at which this occurred was dependent upon the volume fraction of sand. Longitudinal splitting of the outer plies was also observed before ultimate failure. The initial modulus of the OO/Sand/Oo laminate (Eel) neglecting any ply interaction is given by

o i Eel = E 1 b + Es d / (b+d) ... (1)

o i where Es ' El are the Young's moduli of inner sand lamina and glass outer lamina respectively, band d are the outer and semi-inner ply thicknesses. The modulus of inner sand lamina on its own can be evaluated by Law of mixture equations /6/ for limiting conditions of uniform stress and uniform strain. For non-bonded particles it is considered that poth particles and matrix can carry equal stress. The lower bound for E~ is therefore given by:

667

i Es = Es Em/ (Es Vm + Em Vs ) ••• (2)

Where Es ' Em' Vs and Vm are the moduli and volume fractions of sand and matrix respectively. In the case of a strong bond between the individual components an upper bound obtained from consideration of strain uniformity is obtained:

i Es= Es Vs + Em Vm ... (3)

Table 1 shows the comparison of calculated and experimental Young's moduli for the graded sand filled OO/Sand/Oo laminates. The laminate modulus is much higher than expected for a laminate in which the inner sand I ami na \,as cons idered to have 1 imited load bear ing abi 1 i ty because the sand particles had not been treated with an adhesion promoter nor did the literature suggest that sand polyester composites had upper bound moduli. It is seen that when Es takes a lower bound solution (Eg.2) it is a gross underestimate, whereas the upper bound tends to (Eg 3) overestimate the observed laminate modulus. However, a higher modulus to the sand filled resin has been observed /3,7/ where it \,as concluded that the modulus of the sand filled resin resulted from the agglomeration of the irregularly shaped particles and the presence of an heterogenous residual thermal stress which overcame the lack of chemical adhesion between the particles and the matrix, to produ('(' a load bearing interface /3/. Filler agglomeration is l'e~3pOlisibl(' for a modification in the local load transfer mechanism, from shear load bearing to particle-particle load bearing, and increases the load bearing capacity of the composite accordingly. The modulus of filled res ill in lerms of particle shape, size and degree of aggregatioll has lJE'en discussed previously /7/. A better estimation of 18minl1t" modulus is "bt aillt'd by translating the modulus of the sand fillE'd resin coupon when tested independently, into the laminate. The agreem' 'nt between the measured values and calculated modulus from equation 1 are shown in Table 1, to be good. Similar agreement was also found for finp sand filled OO/Sand/Oo laminates (Table 2). Thus, the sand lamina has contributed towards the initial stiffness of the laminate and carrips an approximately equivalent load to a continuous glass fibre laminate. However, the failure of the inner sand lamina in OO/Sand/Oo occurs by microdamage accumulation which is analogous to the transverse failure of 90° ply in 0°/90% ° cross ply laminates. The photomicroscopic st.udies suggest that the cracks nucleate by particle de~etting. Kith thick inner sand layers, transverse cracks propagate l'a[Jidly across the v,idth of the specimen, ~hereas for laminates wi th thin inller pi ie'S, a two stage process was observed. Dev,etting occurring at small strains but propagation being constrained until higher composit" strains v"ere rpached. Constraint at low inner pI:; I h i ckllesses demons t I'ates the analogy with cross ply laminates.

The e'Jlla t ion for the mill i mum t ransl'erse c racking strain according to Bailc~ ct al /8/ for a 0°/90%0 laminate can be used analogously for OO/Sllncl/Oo laminate alld is gil'(,11 by:

668

E~lU (min) = [2 Ys b E~ (lit / E~ Eel (b+d) t- E:~

where (Ii = Eel G! (b+d) / E~ E! b d2

... (4)

The fracture surface energy of sand lamina, Y s' was measured by work of fracture method using SENB specimens and taken to be 55 Jm2 [5]. ~ is the shear modulus of the sand lamina and taken as 10.5 GPa. The magnitude of the thermal strain E}r which developed in the longitudinal direction in the inner sand lamina was calculated by the following equation. th 0 0 i

Esl = El b (as - all (T1 - T2) / (El b + Es d) ... (5) in which (as- all (T1 - T2 ) was determined from the radius of curvature of DO/Sand unbalanced beam /2/.

Fig 3 shows the comparison of transverse cracking strain predicted from equation 4 and the onset of microcracking in inner sand lamina as a function of d. Agreement for both graded and fine sand filled laminate is good, and the equation predicts the observed cracking well. At large d, the inner sand lamina fails at a strain equal to the failure strain of the equivalent sand filled resin (Esu )' The thermal correction when added to the observed inner ply cracking strain gives value close to Esu' At small d, the constraint effect operates in these laminates, and the inner sand lamina fails at a much higher strain than Esu' However, the failure strain of the inner sand lamina tested on its own cannot be represented by a unique value, since the strength of a filled system depends, on the average particle size /9/. With coarse particles, dewetting occurs at relatively low stresses because the particle interfacial area decreases with increase in particle size. In graded sand filled laminates the largest particle reaches a size of 1.18 mm, and is therefore expected to fail at much lower strains in contrast to fine sand filled laminate where the maximum particle size is 600 pm. This is confirmed by relatively lower cracking strains for the graded sand filled laminates (Fig 3). Since, equation 4 neglects the origin of the crack or the way it propagates through the matrix, it is therefore reasonable to conclude that the constraint theory originally developed for 00 /900 /00 cross ply laminates can be applied confidently to the behaviour of DO/Sand/Do glass/sand hybrid laminate.

4 - CONCLUSIONS The stiffness of aO/Sand/Oo laminates can be closely predicted from the a rule of mixtures when the modulus of sand filled resin tested independently is translated into the laminate, although the upper bound gives a better estimation from component properties. It should also be noted that the residual compressive stresses considered responsible for the sand/matrix composite properties are not destroyed by the small thermal stresses induced by lamination. The change of slope in the stress-strain curve is attributed to dewetting of sand particles and simultaneous disruption of the agglomerates. The failure

669

of the inner sand lamina occurs by accumulation of microdamage. The lamina cracking behaviour can be described by a modified cracking constraint theory which demonstrates the integrity of the composite. Once the stress transfer mechanism between the individual laminae are understood, there is potential for extending GRP pipes with particulate fillers.

Acknowledgements We thank the Government of Pakistan for financial assistance, Scott­Bader & Co Ltd for resin and Stanton PIc for supplying the sand.

REFERENCES 1. S.W.Tsui and F.R.Jones Proc. of Int. Conf. on Testing,

Evaluation and Quality control of Composites-TEQC 87 Guildford, UK, Sept 1987.

2. S.Ahmed and F.R.Jones Proc.3rd.lnt.Conf.Fibre reinforced composite'88 Extending the limits {PRI,Liverpool,UK,March {1988} paper 16

3. S.Ahmed and F.R.Jones Composites 19 {1988} 277. 4. F.R.Jones M.Mulheron and J.E.Bailey J.Mater.Sci. 18

{1983} 1522. 5. S.Ahmed and F.R.Jones in preparation. 6. L.J.Broutman and R.H.Krock Modern Composite Materials,

{Addison Wesely, Massachusetts} 1967. 7. S.Ahmed and F.R.Jones submitted to Composites. 8. J.E.Bailey P.T.Curtis and A.Parvizi Proc. Roy. Soc. LOlldon

A366 (1979) 599. 9. H.Hojo and W.Toyoshima 31st ANTEC,SPC,Montreal,Canada

1973, pp163

lii..:; 1. Schematic uiagram of mouel OO/Sand/Oo hybrid glass [ihr'e laminale.

200

I 40

~--..'--r---'.81~21"-o . .----x'

StraIn (%)

Fig 2. Typical sLl'ess-sLrllill CUI've fUL'

00 /Santi/Oo comp' '" i 1.(" hi lIL ,b=O. H lind d=l.l mm, \°[,=17% alLd \J.5 =28X.,

670

Table 1. T.he theoretical and experimental values of the low strain Young's moduli of DO/graded sand/Do (I) and DO/fine sand/Do (II) GRP laminates respectively.

I ITablel d b

(mm) (mm) (X) (X) (OPa) (OPa)

(1) 0.87 0.56 18.30 25.44 19.55 15.20 1. 21 0.63 18.30 29.85 23.53 14.85 1. 28 0.96 16.82 29.76 18.63 16.25 2.2D5 0.77 18.30 30.59 22.74 12.53 2.655 0.62 19.13 30.59 18.87 10.85 2.605 0.70 18.30 30.98 20.47 10.59

( 2: ) 0.72 0.80 18.67 26.08 23.95 19.95 0.72 0.70 17.62 17.55 20.54 17.85 O.Bl 1. 12 19.93 19.06 22.64 12.94 1. 79 1. 02 20.01 29.32 21. 65 16.14 1.92 0.93 29.04 29.04 22.05 14.91 2.165 0.70 21. 22 28.38 17.26 12.98 2.595 I. 44 22.43 30.03 16.46 13.00

e lb ub m

(OPa)

22.30 24.98 23.08 21. 35 20.49 20.23

24.59 19.47 25.05 22.46 21. 61 20.31 23.52

ub Ed

(GPa)

25.76 27.20 27.02 26.87 26.42 26.42

27 48 23.83 27.38 24.93 26.65 26.70 27.27

Ecl is the experimental modulus, Ecl' Ecl' Ecl predicted values from equation 1, where ~s is either the calculated values from law of mixture equations lower bound (lb) and upper bound (ub) or the measured modulus (m) of a sand filled resin coupon respectively/3/.

0.2

0.16

~

:e ~ 0.12 -

H rJIl

0.08 -

0.041,,-0--.,..1

Fig 3. Comparison of the theoretical and experimental values of inner lamina cracking strain Eslu as a function of semi-inner ply thickness d for DO/Sand/Do laminate.--- predicted curve from equation 4 using b=l, G~=ID.5 GPa, V~=26% and Vf=18% .• DO/Fine sand/Do, ~ DO/Graded sand/Do laminate.

COMPRESSIVE BEHAVIOUR OF UNIDIRECTIONAL GLASS/CARBON HYBRID LAMINATES

ABSTRACT

G. KRETSIS, F. MATIHEWS, J. MORTON·, G. DAVIES

Imperial college - Department of -1eronautics Prince Consort Road - SW7 2BY LONDON - England

'Virginia Poly Institute & State University E.S.M. Department College of Engineering

BLACKSBURG Virginia 24061 - USA

Experimental data on the compressive strength of glass carbon hybrids are presented. Carbonlglass ratios of 0, 12.5, 25, 37.5, 50, 75 and 100% were tested in 1 and 2mm thick unidirectional laminates. It was found that lower strength was obtained with carbon plies on the surface of the specimen. a-ply were found to be stronger than 1S-ply specimens, a result confirmed by a finite element investigation of the specimen stress distribution. The rule of mixtures was found to give a good estimate for hybrid modulus, provided the stiffness of the constituents was taken at the appropriate strain level.

INTRODUCTION

Composites containing more than one type of fibre are commonly known as 'hybrid composites'. The level of mixing can be either on a small scale (fibres, tows) or a large scale (layers, ribs). The purpose of hybridisation is to construct a new material that will retain the advantages, but not the disadvantages, of its constituents. For most properties the rule of mixtures, (Le the volume-weighted sum of the constituents' properties), proves to be only an upper bound. The current paper presents one aspect of the results obtained from a large experimental and theoretical programme in which both unidirectional and multi-directional laminates were studied. Other parts of the programme have been reported elsewhere 11,2,3/.

- MATERIALS

Unidirectional hybrid composite laminates were fabricated by laminating pre-impregnated sheets of either XAS carbon or E-glass fibres in Ciba-Geigy 913 epoxy resin. Both a-ply and 1S-ply symmetric laminates were produced, of nominal thickness 1 and 2mm respectively at SO% fibre volume. Hybrid mixtures from all-glass to all-carbon were studied, the full range of stacking sequences being given in Table 1.

672

2 - EXPERIMENTAL PROCEDURE

After cutting from the laminates, specimens were left in the appropriate environment for conditioning. Those to be tested 'dry' (the majority) were kept in a sealed dessicator over silica gel for 4 months. Those to be tested 'wet' were kept sealed, at 85% relative humidity, over saturated potassium chloride solution at room temperature. The time available only permitted about 90% of full saturation to be reached. It was found that the carbon-rich specimens absorbed more moisture than the glass-rich laminates. Uniform thickness end plates, of aluminium alloy, were glued to the specimens after conditioning, the specimens then being returned to the appropriate environment for a short time before testing.

Specimen geometry conformed to the CRAG specifications 14/. The gauge length-to-thickness ratio was 5:1 for all specimens, and the width was held constant at 10mm. The specimens were mounted in a Celanese jig and loaded in a Zwick displacement controlled machine. In generaliS specimens per configuration were tested. Strains were monitored by strain gauges affixed to both faces of a number of specimens.

3 - EXPERIMENTAL RESULTS

Most of the specimens failed at the junction of the gauge length and the end plates. A few specimens failed centrally and a few under the end plates; the latter were regarded as invalid. Microscopic examination showed evidence of fibre micro-buckling, especially in those specimens failing at the end of the end plates. Extensive shear cracking was also observed.

Strengths seemed to be unrelated to failure modes and the scatter was higher than for tension tests 11/. Saturated specimens gave fewer valid, and weaker, failures than dry specimens. Details of the strengths, together with some comparative tensile data, are given in Table 2. It was found from the strain measurements that, for the all-carbon and hybrid specimens, compressive modulus decreased with increasing strain; this was thought to be caused by micro-buckling of the carbon fibres.

It is common to express the performance of hybrids in terms of the 'hybrid effect' (the enhancement in failure strain of the low elongation phase (carbon, here) when part of a hybrid). This can be determined in tension because the carbon fibres will fail before the hybrid reaches its ultimate strain. In the compression tests reported here it was not possible to distinguish initial failure from ultimate failure and hence a hybrid effect could not be determined.

4 - THEORETICAL TREATMENT

Very little theoretical work is available for predicting the compressive performance of hybrids 151. One major difficulty is the sensitivity of strength measurements to the test method employed. In addition none of the traditional models for strength prediction, which all overestimate strength, are suitable for hybrids. In the current programme most of the theoretical effort was expended on the tensile and flexural aspects of the work.

It was found here that the rule of mixtures was an acceptable method for predicting the stiffness of hybrids, provided the moduli of the constituent phases were taken at the appropriate strain level.

673

A limited two-dimensional finite element study was undertaken to look at the influence of the end plates on the stress distribution in the specimens. Some agreement with experiment was established. For the all-carbon lay-ups the 8-ply specimen was predicted to be 28% stronger than that with 16-plies; experimental results showed only an 18% difference. However the model ignored thermal strains, statistical effects, three-dimensional effects, etc. It was also shown that for carbon-rich hybrid lay-ups a length-to-thickness ratio of 5:1 was not sufficient to give a uniform through-thickness stress field at the centre of the gauge length.

5 - DISCUSSION

Compression tests are notoriously difficult to carry out successfully, the results, particularly for strength, being very sensitive to test method. In the current programme an early decision was made to conform to the CRAG specifications 14/, which for compression implied using the Celanese jig. This method gives similar results to those obtained from using the IITRI jig, although other methods can give higher strength 16/. It is believed that the method of loading, which introduces shear along the end plates, causes axial strain concentrations at the end of the end plates. This is confirmed by the finite element studies and the experimental results; most specimens failing at this point or under the end plates.

Although failures under the end plates were taken as invalid it was, in fact, difficult to distinguish between the strengths associated with the different failure modes. It was felt that these failures were caused by uneven gripping, brought about by either non-uniform thickness or poor alignment of the grips of the loading jig. The moisture conditioned specimens gave fewer valid failures than the dry specimens, problably due to the deleterious effect of moisture on the adhesive used to bond the end plates.

The reduction of compressive modulus with strain for the all-carbon and carbon-rich hybrid specimens was thought to be caused by fibre micro-buckling. This is consistent with the all-glass specimens (larger diameter fibres) showing no change of modulus with strain. Certainly, there was considerable evidence of fibre micro-buckling seen in the micrographs of failed specimens.

6 - CONCWSlONS

From the present study the following conclusions can be drawn:-Compression strength is reduced for specimens with carbon on the outside (as is tensile strength). Moisture reduced compression strength only when carbon was on the outside of the specimen. 8-ply are stronger than 16-ply specimens. For all-carbon specimens compression strength is less than tensile strength; the opposite is true for all-glass. The compression strength is approximately the same for all-carbon and all-glass lay-ups. Stiffness decreases with strain for hybrid and all-carbon lay-ups. Stiffness is adequately predicted by the rule of mixtures, provided constituent modulus values are taken at the appropriate strain level.

7 - ACKNOWLEDGEMENTS

This work was carried out with the support of the Procurement Executive, Ministry of Defence, UK.

674

8 - REFEROCES

- Kretsis G., Matthews F.L., Morton J. and Davies G.A.O., Proc. Cont. "Fibre Reinforced Composites 86", Liverpool, UK, (April 1986) 225-229, Institution of Mechanical Engineers, London.

2 - Kretsis G., Matthews F.L., Morton J. and Davies G.A.O., Proc. Cont. "ICCM-VI/ECCM-2", London, UK, (July 1987), Vol 1, 221-230, Elsevier Applied Science, London.

3 - Kretsis G., Matthews F.L., Morton J. and Davies GAO., Proc. Conf. "Engineering Applications of New Composites", Patras, Greece, (August 1986), 421-432, Omega Scientific, Wallingford, UK.

4 - Curtis P.T., Technical Report TR 88012, (February 1988), Royal Aerospace Establishment, Farnborough, UK.

5 - Kretsis G., Composites, 18 (1987) 13-23. 6 - Haberle J.G. and Matthews F.L., "The compressive mechanical properties of

fibre-reinforced plastics", Progress Report (November 1988) MoD(PE) Agreement No 2037/325 XAIMAT, Dept. of Aeronautics, Imperial College, London, UK.

Table I StackIng Sequences Tested Single layer thIckness = 0.125mm when cured (nomInal) FIbre volume fraction = 0.60 when cured (nomInal) D = dry, S = saturated. All lay-ups symmetrIc.

StackIng Sequence Number

6 46

7 8 9

10 11 12 13 14 15 16 17 18 19 20 21 22 23 24

3 4

StackIng Sequence c=carbon g=glass

mId-thIckness-+( [gggg]

[gggggggg] [gcgggggg]

[cggg] [ccgggggg]

[gggc] [gcgg]

[gccggggg] [gccgcggg]

[ccgg) [ccccgggg)

[ggcc) [cgcg] [gccg)

[gcgcgcgc] [gccgccgg)

[cccg] [ccccccgg)

[gccc] [ggcccccc]

[ecce] [cccccccc]

Carbon Vol %

o o

12.5 25 25 25 25 25 37.5 50 SO 50 SO 50 SO SO 75 75 75 75

100 100

CondItIon

D D D D D D D D D D D D D D D D D D D D D D

s

s s

s s

s s s S

s s s

S S

675

Iabl~ ~, I~Sl B~sullS

- DRY - - WET-

Stack cr(MPa) crcompr E(%) Ecompr cr(MPa) Seq. ult. ±COY crtens ult. gens ult. ±COy

6 1432 ±10.5% 1.48 1289 ±14.4% 46 1475 ±9.9% 1.42 3.00 1.24

7 1051 ±3.8% 1.14 1.77 1.13 8 1171 ±7.2% 1.15 1083 ±7.6% 9 942 ±5.8% 0.87 1.50 1.02 883 ±3.7%

10 1452 ±6.1% 1.30 11 1479 ±7.3% 1.32 1352 ±14.4% 12 1053 ±12.9% 0.90 1.66 1.06 1048 ±10.1% 13 1074 ±13.0% 0.82 1.47 0.96 14 1376 ±4.3% 1.01 1258 ±7.4% 15 1150 ±7.4% 0.75 1.36 0.91 1085 ±6.0% 16 1607 ±8.4% 1.13 1589 ±10.9% 17 1211 ±10.0% 0.81 1298 ±9.0% 18 1569 ±7.2% 1.07 19 1284 0.83 1.55 1.01 20 1274 ±12.1 % 0.85 1.54 1.03 21 1549 ±8.4% 0.89 1465 ±9.0% 22 1118 ±7.8% 0.63 1.00 0.70 1196 ±8.1% 23 1676 ±6.4% 0.91 1673 ±6.3% 24 1359 ±6.4% 0.74 1.30 0.92

3 1628 ±7.4% 0.76 1492 ±14.6% 4 1385 ±6.9% 0.62 1.09 0.79 1202 ±7.3%

coy coefficient of variation

FATIGUE OF HYBRID COMPOSITES

B. HARRIS, T. ADAM, H. REITER

University of Bath - BA2 7 A Y BA TH - England

ABSTRACT

A study has been carried out of the fatigue behaviour of hybrid composites of carbon/Kevlar-49 and carbon/glass in epoxy resin. For these two families of materials a comparison is presented of the prop­erties of unidirectional composites and laminates with a [(±45,02)2]S structure. The unidirectional composites have been tested in repeated tension and tension/compression at R ratios between +0.1 and -1.2. All other materials have been tested in repeated tension fatigue only (R = +0.1). The results have been analyzed by normalizing relative to the elastic modulus and to the ordinary tensile strength, revealing a homo­geneous pattern of behaviour for both families of hybrids.

I - INTRODUCTION

Aramid fibre composites exhibit an unusual type of fatigue response with a steeply descending stress/life characteristic and drastic shortening of the fatigue life at peak stress levels below about 90% of the normal failure stress[1]. A potential consequence of this was that in forming a carbon/Kevlar hybrid, the apparent weakness of the Aramid fibre might detract unduly from the otherwise excellent fatigue response of the CFRP. A second possibility was that a loading regime with a compression component could lead to an even worse performance of the KFRP constituent in a C/K hybrid. We have therefore carried out a study of the fatigue behaviour of two families of hybrid compos­ites in order to compare their response with that of single fibre composites.

II - MATERIALS AND TESTING PROCEDURES

The materials studied were of two kinds, carbon/Kevlar-49 (C/K) hybrids with a Ciba-Geigy 914 epoxy matrix, and carbon/glass (C/G)

678

hybrids with a 913 epoxy matrix. The carbon fibres in each case were Courtauld XAS fibres. Experiments were carried out on two families of unidirectional hybrid laminates, with a range of compositions between CFRP and KFRP on the one hand and CFRP and GRP on the other. For comparison some experiments were also carried out on plain CFRP, plain GRP, plain KFRP, and 50/50 C/K and C/G hybrids with a [(±45,Oz)2]s layup. The stacking sequence for these laminates was [(±45)X,(0)C,(0)X,(±45)C,(0)X,(0)C]s, where X is either Kevlar-49 or E­glass. After autoclaving, the composites were conditioned to a mois­ture content of approximately 1%. Test samples were straight-sided, 20mm wide and 2mm thick, with a free test length of 100mm between aluminium end tabs. The unidirectional C/K samples were fatigued under constant load at R ratios, Omra/Oux, between +0.1 and -1.2, anti-buckling guides being used for tests involving compression. The majority of fatigue tests on the unidirectional C/G composites and on the [(±45,Oz)z]s laminates were carried out at an R ratio of +0.1 (ie without a compression component). Fatigue tests were carried out under load control at a rate of load application of 200kNs-1.

III - EXPERIMENTAL RESULTS

1. Unidirectional Composites

The composition dependences of the strength, failure strain and elastic modulus of the two families of unidirectional hybrids are shown in fig 1. The results in this figure are entirely predictable from simple mechanics (the moduli) and the simple failure strain modeI[2] of hybrid failure (the strengths).

Stress/log life (S/logNf) curves for the two sets of hybrids in repeated tension cycling (R = +0.1) are shown in fig 2 (the original data points have been omitted to improve clarity). The mixed fibre composites appear to behave in a uniform fashion dictated in a rela­tively simple way by the proportions of the two fibres present. The overlapping pattern of the curves for the C/G family is initially a consequence of the non-linear variation of strength with composition. For the C/K composites, the effect of the compression stress component can be seen in the constant life diagram for an endurance of 105

cycles (fig 3). This shows a homogeneous response through the compo­sition range. The peak stress for a life of 105 or 106 cycles is a linear function of composition regardless of the level of compressive stress in the cycle[3]. The presence of the Aramid fibre exerts no damaging effect on the fatigue response of the hybrid and, contrary to expecta­tion, the susceptibility of KFRP to weakness under compressive loads appears to be no greater in fatigue than under monotonic loading conditions. By contrast, the results for the C/G composites in fig 2 lead to a peak stress vs composition relationship with a distinct positive deviation from linearity, suggesting that hybridizing the CFRP with GRP may result in a more positive benefit than incorporating KFRP.

The stress/life curves of fig 2 are replotted in fig 4 to show fatigue life as a function of initial maximum strain for repeated tension cycling (data points now included). This is effectively a normalization

679

relative to the elastic modulus, by contrast with the fatigue ratio approach. Despite the inherent scatter in all of these fatigue tests, there is a similarity of patterns of behaviour within each of the two hybrid families. In each family of curves most of the experimental points fall within a single scatter band except for those of the second single fibre species (ie the plain GRP and KFRP). This aspect of the behaviour of the unidirectional KFRP composites was also found to be retained at R values other than +0.1[3]. The pattern is so definite in each case that it seems likely that the scatter may simply be a conse­quence of the combination of compositional variability and the inher­ently stochastic nature of the fatigue process itself.

2. [(±45,0z)z]s Laminates

Table 1. Properties of [(±45,02)Z]S CFRP/KFRP and CFRP/GRP Laminates.

Composite Young's Tensile Failure Modulus Strength Strain

GPa GPa %

914/XAS carbon 81.1 1.28 1.39 CFRP/KFRP(50:50) 52.7 (0.92)* 0.81 (0.84) 1.51 (0.94) 914/Kevlar 33.1 0.64 1.82

913/XAS carbon 76.0 1.17 1.47 CFRP/GRP(50:50) 51.7 (1.03) 0.82 (0.87) 1.82 (1.02) 913/Glass 23.9 0.71** 2.11

I hgures In brackets are the ratios of the actual hybnd properties to the leans of those of the single fibre cOlposites, eg lEH!(EclEK).

IImsured at the sale testing rate as the fatigue tests.

The S/logNf curves for these two groups of materials were similar in shape and disposition to those of the related unidirectional compos­ites given in fig 2. When the results are presented in terms of initial strain versus life (fig 5) it can be seen that differences in strength and stiffness resulting from the incorporation of the 45' plies do not, for the most part, change the fatigue behaviour of the hybrids of either family. The three curves for the C/G materials are exactly superposable on the equivalent curves for the unidirectional compos­ites shown in fig 3. In the case of C/K, the results for the plain CFRP and the 50:50 hybrid results are also for the most part identical to those for the corresponding ud composites, but the points for the plain KFRP laminate fall above the band of points from the other two compositions. This does not occur in the unidirectional materials and is the only significant difference between the sets of unidirectional and [(±45,02)2]S pairs in each family.

IV - DISCUSSION

On a strain/life basis there are marked similarities in behaviour between the unidirectional composites of the two hybrid families and the corresponding [(±45,Oz)z]s laminates, fatigue response, for the most

680

part, being reasonably accurately predicted for almost any composition, of either layup, from a knowledge of the strain/life curve for any member of the family.

If the fatigue responses of the two hybrid families are compared by normalizing the constant life data in terms of the monotonic tensile failure stress, ie. in terms of the conventional fatigue ratio, the two families of hybrids show almost identical response (fig 6), and in both cases there is a positive deviation from linearity. Thus, the factors which determine monotonic strengths are not simply carried over into cyclic loading response. We note also that the fatigue ratios (106

cycles) for the two 50:50 [(±45,02)2]S laminates are 0.72 (C/K) and 0.73 (C/G), almost identical with the two values of 0.75 for the unidirec­tional materials of 50:50 composition in fig 6, and show approximately the same level of "synergism". It is of course not possible to describe this as a synergistic effect, since we have no detailed mechanistic model of hybrid fatigue damage accumulation against which to assess the actual result.

v - CONCLUSIONS

.Studies of the fatigue behaviour of unidirectional and [(±45,02)2]S laminates of two families of hybrid composites have given conventional stress/life data that appear to fall into relatively simple patterns in relation to composition •

• Two kinds of normalization procedures have been used in order to compare the behaviour of the two kinds of hybrids and of the two different structures. The first is to plot strain/life curves (normalization relative to elastic modulus) which provides data that are more relevant to a designer, and the second is to normalize relative to tensile strength, which is the conventional way of obtaining the fatigue ratio. The two methods emphasize slightly different aspects of behaviour, but by and large both give indications that hybridization of CFRP with either an Aramid fibre or with glass fibre offers potential benefits in respect of fatigue response that are not in any way pre­dicted by the ordinary mechanical properties of the hybrid composites •

• The results suggest that the problems of designing with hybrids for a fatigue environment may be somewhat simpler than might otherwise be expected, since interpolation appears to carry no hidden dangers.

ACKNOWLEDGEMENTS

The detailed experimental results on which this analysis is based are to be published in due course[3,4,5]. The authors are grateful to the Procurement Executive (MoD) for sponsorship of this work and to Dr PT Curtis and Dr G Dorey of RAE, Farnborough, for their advice and interest.

REFERENCES

1. Jones CJ, Dickson RF, Adam T, Reiter H and Harris B, Proc Roy Soc Lond, A318 (1984) 461-475.

681

2. Chou T-W and Kelly A, Annual Reviews of Materials Science, (1980) .10. 229-259.

3. Fernando G, Dickson RF, Adam T, Reiter H and Harris B, (1988), Fatigue Behaviour of HTbrid Composites: I Carbon-Kevlar HTbrids, J Materials Science, in press.

4. Dickson RF, Fernando G, Adam T, Reiter H and Harris B, (1988), Fatigue Behaviour of HTbrid Composites: II Carbon/Glass HTbrids, J Materials Science, in press.

5. Adam T, Fernando G, Dickson RF, Reiter H and Harris B, (1988), Fatigue LUe Prediction for HTbrid Composites, International Journal of Fatigue, in press.

"' Cl. l:J

i5 b~

2

o I o

I 10

I 40

I 60

I 80

Vol % CFRP

I 100

Fig.! Co.position dependence of .echanical properties for unidirectional elK and e/G hybrids.

2 2

~ ~KfRP

lS".KAlP ~

SO,..KFRP ~ R= .0·' KfRP

~ 0 0 2 4

R=.Q·, Log Nr o ~~ ____ ~ ____ L-__ ~~ __ ~

o 2 4 6 8

Leg Nt

Fig.2 S/iogHf curves for unidirectional hybrids; a) elK, b) C/G (R : to.1).

GRP

6 8

682

ro a... L:J

~ 0

2

0

1.5

1.0

05

0 -1.5 -to -Q5 0 0.5 1.0 2.0

Fig.3 Goodlan diagraa showing effect of R ratio and cOlposition on fatigue of unidirectional elK cOlposites

• 0 __

;, -;---~, -_ 0"'.

-.pe__ _ a~O ~ • • -If ......

° iC !l,"

o " • (fRP o • lS,..KfRP

· 36% o 0 SO" 0 KfRP

0 2 4

Fig.4 S/logHf curves of fig 2 replotted as strain/log life curves (R : 10.1).

GRP

3

2 • • • ••

• GRP • 15%GRP .. SOo/a

• 25% • CFRP

Log Nt

I £"'ax % •• 2

•••• • .{ .~

• • • CfRP .. • KFRP

0 • (fRP/KFRP 0 0 2 6 0 2 4 6

Log N, Log N

Fig.S Strain/log life curves for cOlposites of [(±45,02)2)S construction (R : 10 .1).

683

vr 10 Q)

u >-w CFRP/GRP

-.0 C> \ --~ --..", 0 ,-

-+- ./ ro // \FRP/KFRP '-Q) :::J 0'1

R=+0'1 :;: ro

LL 0

0 20 40 60 80 100 Vol % CFRP

Fig.6 fatigue ratio (106 cycles) as a function of co.position for elK and CIG unidirectional co.posites (R : 10.1).

MODELISATION, SINGULARITES MODELING

Chairman: Pro J. C. SEFERIS University of Washington

AN ANALYTICAL INVESTIGATION ON THE THERMALLY INDUCED RESPONSE OF COMPOSITES

ABSTRACT

IN THE ABSENCE OF THERMAL EQUILIBRIUM

J. FLORIO Jr, J.B. HENDERSON, F.l. TEST

University of Rhode Island - Department of Mechanical Engineering and Applied Mechanics Wales Hall 02881-0805 KINGSTON, R.I. - USA

A numerical study of the thermally-induced response of two widely used composites has been conducted. This study was performed using a newly developed numerical model which, unlike previous models, does not include the assumption of local thermal equilibrium between the solid matrix and gases generated as a result of the decomposition processes. The results of this study include temperature, mass loss, pressure and expansion profiles. These results were used to evaluate the effects of material composition and processing on the overall response of the materials.

INTRODUCTION

Polymer composites are currently used in a wide variety of thermal protection applications. Their overwhelming acceptance in thermal protection systems is primarily due to the thermal and transport properties of the carbonaceous char matrix which is formed as a result of the low temperature pyrolysis reactions. To efficiently design thermal protection systems, it is necessary to be able to predict the thermally-induced response of these materials a priori.

The purpose of this work was to study the thermally-induced response of two similar glass filled composites, using a newly developed numerical model. Unlike previous work /1,2/, this model includes the assumption of local non-thermal equilibrium between the solid material and volatiles, thermochemical expansion, and related work terms incorporated into the energy equations.

I - MATERIAL BEHAVIOR

When a polymer composite is exposed to a surface heat flux, the initial heat transfer is primarily due to transient energy conduction. When this material reaches sufficiently high temperatures (200 - 300·C), chemical reactions begin to occur. These thermally-induced reactions, commonly referred to as pyrolysis reactions, result in the degradation of the resin component of the composite matrix into residual char and product volatiles. This reaction zone moves from the heated surface through the

688

material. During the initial stages of the pyrolysis reactions, decomposition gases are trapped with the pore network due to low material permeability. This accumulation of gases results in the internal pressurization of the material. For many polymer composites the porosity and permeability are small enough to cause the internal pressure to become quite high. This internal pressurization is at least partially responsible for a very rapid and sometimes large expansion of the material which occurs at this time. As a result of the continuing expansion and decomposition the material permeability and porosity begin to increase. This increased permeability, coupled with exiting pressure gradients, results in the flow of gas through the pore network. The gases which flow back through the char structure remove energy, thus attenuating the conduction of heat to the reaction zone. The gases which flow through the virgin material serve to pre-heat the material. The rate at which energy is transferred between the solid and gas in the pore network is characterized by what is known as the volumetric heat transfer coefficient, which in general is dependent on the stage of decomposition and the flow rate of gas. As the pyrolysis reactions proceed, the permeability and porosity increase still further. This results in increased gas flow and in a reduction of internal pressure. When this occurs the material experiences a rapid contraction. At temperatures in excess of !OOO·C, the carbonaceous char reacts with the silica, present in glass reinforced composites, causing additional changes in the thermal, transport and mechanical properties of the material.

2 - MODEL DESCRIPTION

A one-dimensional transient mathematical model has been developed to predict the thermally-induced response of polymer composites without the assumption of local thermal equilibrium. The introduction of convective thermal transport between the solid and product gas requires the solution of a conservation of energy equation for each phase. These equations are coupled through the volumetric heat transfer coefficient, which characterizes the rate of energy transfer between the two phases within the pore network. The mathematical model also includes appropriate equations which describe the following physical processes: conservation of mass and momentum; solid decomposi­tion; thermochemical expansion; and lastly, variable material and transport properties. For the sake of brevity, these equations are omitted from the text. However, they are presented in detail elsewhere /2,3,4/. In the numerical scheme, these equations were solved using a fully-implicit finite difference technique.

3 - MA TERIAL DESCRIPTION

The two composites studied are designated as H4lN and MXBE-350, and are fabricated by Ametek, Haveg Division and Fiberite Corporation, respectively. Both materials consist of a basic phenol-formaldehyde (phenolic) resin with varying quantities and types of fillers and reinforcing agents. These materials were chosen for study because they display typical decomposition/expansion behavior for glass-filled composites, and are widely used in thermal protection applications.

Haveg H41N consists of a phenolic resin containing short glass fibers and talc. The manufacturing procedure consists of mixing the filler components into the resin such that the glass fibers are randomly orientated. However, as a result of the processing the glass fibers become somewhat orientated in the direction normal to the applied pressure.

Fiberite MXBE-350 consists of an acrylonitrile-butadiene (rubber) modified phenolic resin containing glass fibers and glass powder. However, unlike H4IN, the glass fibers in MXBE-350 are in the form of woven glass mats, making MXBE-350 a laminated material (layered). As a result of their respective glass fiber orientations, both materials are somewhat transversely isotropic. For example, both materials expand significantly in the direction normal to the preferred fiber orientation, while displaying little expansion in other directions. The material and transport properties for H4lN and MXBE-350 have been presented elsewhere /2,3,4/.

689

4 - RESULTS AND DISCUSSION

Using the appropriate equations and thermal property relations outlined earlier, the thermally-induced responses of H41N and MXBE-350 were evaluated. The model was exercised for a .03m thick slab with a surface heat flux of 279.7 kW /m2 applied to the front surface, x = O. This heat flux provides sufficient thermal energy for the initiation of only the pyrolysis reactions. The carbon-silica reactions were not investigated in this study. The back surface was prescribed to be insulated. The initial temperature and pressure were taken as 40'C and 1.01325 x 105 Pa, respectively.

For both materials the model was exercised for a volumetric heat transfer coefficient of 2.0 x 105 W /m3K. This value was experimentally determined from preliminary heat transfer coefficient measurements of H41N.

The thermally-induced response of H41N and MXBE-350, as a function of the time and depth are shown in figures 1-4 and figures 5-8, respectively. All profiles are shown at times of 100, 200, 400, 600 and 800 s.

Figure I illustrates the temperature histories of the solid material and product gas for H41N. As can be seen, the major deviation from local thermal equilibrium occurs at the heated surface. This is expected because this surface is exposed to the heat flux, causing rapid thermal changes in the solid material. The volumetric heat transfer coefficient is too small to induce thermal equilibrium at or near the surface, especially at the early times when the net absorbed energy at the surface is highest. However, as time increases the gas temperature at this surface approaches the solid temperature, primarily due to the balancing of the absorbed and emitted energy from the solid at the surface. Also evident from this figure are the steep temperature profiles, present primarily because of the low solid thermal conductivity. The expansion of the material from an initial thickness of .03m to approximately .032m at 800 s is also seen.

Figure 2 shows the mass loss history for H4IN. The ordinate axis represents the fraction of mass remaining. A comparison of figures I and 2 shows that the pyrolysis reactions are initiated in the temperature range of approximately 350 to 375'C, and reach completion in the range of approximately 1000 to 1100'e. As a result, it is evident from this figure that the reactions have reached completion at the heated surface for times greater than 400 s.

The dimensionless pressure profiles for H41N are shown in figure 3. As can be seen, the peak pressure moves further into the material with increasing time. A comparison of figures 2 and 3 shows that for a given time the peak pressure takes place just after the onset of pyrolysis (i.e., ms/ms 0>.98), primarily due to the low material permeability and porosity in this region. The sole exception to this is at 800 s, where it is evident from figure 2 that the pyrolysis reaction zone has progressed through the material, causing a high permeability and porosity throughout. As a result, gas is allowed to flow out the back surface, reducing the entire pressure field.

Figure 4 depicts the fractional length change profiles for H41N. As with the peak pressure, the peak expansion moves further into the material with increasing time. However, examination of figures 3 and 4 shows that the peak expansion lags the peak pressure for a given time. The peak expansion actually occurs when the pyrolysis reactions are approximately 60% complete. This is thought to occur because the char material in the pyrolysis region has a much lower specific strength than the virgin material. Hence, the char structure is weaker, allowing for more expansion as a result of the internal pressurization. This figure also illustrates that the magnitude of the expansion peak decreases with depth due to the decrease in heating rate with depth. This heating rate dependent expansion behavior is discussed in detail elsewhere /4/.

690

Figure 5 illustrates the temperature histories of the solid material and product gas for MXBE-350. A comparison of figures I and 5 shows that while the thermal response of H4lN and MXBE-350 are quite similar, the temperature gradients for MXBE-350 are greater for a given temperature. This is primarily due to the lower thermal conductivity of both the virgin and char components of MXBE-350 at a given temperature. Also evident are the higher surface temperatures associated with MXBE-350, primarily due to the higher surface absorptivity of MXBE-350, coupled with its lower thermal conductivity. This net expansion of MXBE-350 is also seen to be considerably less than that of H4IN.

Figure 6 illustrates the mass loss history of MXBE-350. Inspection of figures 5 and 6 reveals that the pyrolysis reactions begin in the temperature range of approxi­mately 250 to 300°C, and reach completion in the range of approximately 700 to 750°C. Hence, the pyrolysis reactions in MXBE-350 result in a solid mass loss of 30 percent and cover a temperature range of about 500°C. In contrast, the pyrolysis reactions in H4lN occur over a temperature range of approximately 750°C and produce a solid mass loss of 20 percent. This difference in kinetic behavior is probably due to the addition of acrylonitrile-butadiene to the basic phenolic system of MXBE-350.

Figure 7 depicts the dimensionless pressure profiles for MXBE-350. A comparison of figures 6 and 7 illustrates that the pressure dependence on decomposition for MXBE-350 is similar to that of H4IN. However, a comparison of figures 3 and 7 reveals that both the shape and magnitude of the pressure profiles in the two materials are vastly different. As can be seen, the peak value of plpo is approximately 2.3 at 600 s for MXBE-350, as compared to about 12.5 for H4IN. The reason for this is that the permeability of MXBE-350 is approximately 2 orders of magnitude greater than that of H4lN for all stages of decomposition. The high permeability of MXBE-350 most likely is a result of the presence of continuous strand glass mat used in the manufacturing of the material, coupled with diminishing adhesion between the resin and glass mat with decomposition.

Figure 8 depicts the fractional length change profiles for MXBE-350. A comparison of figures 4 and 8 shows that the general expansion behavior of MXBE-350 is similar to that of H4IN. However, the magnitude of the expansion for MXBE-350 is seen to be considerably less than that of H4IN. This is best explained by the lower internal pressures observed in MXBE-350. The relationship between internal pressuri­zation and expansion was discussed earlier. Finally, it can be seen that the expansion of MXBE-350 does not exhibit the heating rate dependence as seen in H4IN.

5 - CONCLUDING REMARKS

The newly developed numerical model presented in this study predicts the thermally-induced response of decomposing, expanding polymer composites without the assumption of local thermal equilibrium. The thermal response of two composite materials has been quantified for a given set of boundary conditions. It has been demonstrated that polymer composites exhibit complex thermally-induced behavior when heated to high temperatures. Also demonstrated is the fact that even seemingly small differences in materials and their method of fabrication have a dramatic effect upon the observed thermal behavior.

ACKNOWLEDGEMENT

The authors wish to acknowledge the financial support of the U.S. Army Research Office, Metallurgy and Materials Science Division, under grants number DAA629-84-K-0192 and DAAL03-86-6-0072.

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A TENT A TIVE INTERPRETATION OF THE CFRP MECHANICAL CHARACTERISTICS BASED

UPON THE FIBREIMATRIX RELATIONS AT THE INTERFACE. A CASE STUDY

J.P. FAVRE, G. DESARMOT, J. HOGNAT", J. ROUCHON"

Onera - BP 72 - 92322 CHA TIL LON - France 'Aerospatiale - Centre commun de recherches Louis Blenot

Laboratoire Central - 12 rue Pasteur - BP 76 - 92152 SURESNES Cedex - France "Centre d'Essais Aeronautiques de Toulouse

23 Avenue Henri Guilaumet - 31056 TOULOUSE Cedex - France

ABSTRACT

Materials gpJectjoo aimed at mmnrat ~ caDs f« rdms cI "efficM!Iq facta- and liOCe/reSin matcbiqr that are bIIed maiDly up<Il a <XIDpIIim cI the (Xl'lS­

tttnents dmacteri!i1ics wIth ~ of the ~ moulded ampaites. Its mown in the p&!IltaIioo UBta full undeiY;nIiCcl tile abageDmJs<3D bed'Aaioed ~ a better knoWledge cI tile fib'e/matrix~ at tile interfaoeo. PJramp1e is taken cI two widely used carton fibres tt300 and AS4) asaodatBd \II1th tile (b 914 resin.

INTRODUCTION

1be~indUstty jSrrequeotty~ wIth the}rotan of se1ec:tq the best fitI'e/resin ~ UBt \G }r0Vide the CCIDpCBte sbU:lUreto bedeslgDec1 wIth the roo;t satistaia 'f ~ cI dmactaistics. 1be usual ~ istogatber asmmyiDf<lJllil.tioos as pcmb1e00 fibres, resiosand~1andnates because tbere is an eridentlact. ct 'ke7S-to selectsysb!ms 00 tile mere basis d tile <mBIitnenIs pq>erties and tile mate­rial ~ used toeaD f<l1h tile "effDD:y fad'«" (ratiod o:mpclIite pt'pIfties to the <XDti.tuents pt'pIfties)and,1M'e lrady, the bazy rmm ('( "litn/r4I6iD ~ (aD ratios taken as a \oIt'tdel

1be ptlIJXa ct the Jresent p!pef is to ~ how a better knoWledge ct the reiati<m between fibres and lDiltrixat tile iDtafaLe <3D <mtritMe to ~ to tile efficiet'q fact« and tilefib'e/matrix~someeJlPelimenta1~.

694

1- BASIC DATA

The owort1mitf to get ateosive data ()Q both ooostibleflts and compoEIibIs was offered r~ /1/ : two carton fibres ttoray T300 and Hercules AS-4) were "aoeaed" wit2l three resins. Fibres, resins and 1arninat.e6 ware tb:lr~ tM»d in varklus CIOOdi­tioos in ceder oot to 1et uodarified any decisive paramet« and to maD <m"eIatioos stand out

())1y two (1 the &yStems eBDined ~ the p-~ namely T300/914 and ASlt/914, are dB:usaed in thef~ : they differ from one ardler oo1y in the fibre as the same be.tdl ct resin was used thr~ the study for plain matrix~, pre­p-egs,1aminates and il:Jterfaoe tssts.

1.1- Basic data ()Q ooostibl9Dts

In agreement W112l mannfacttJrer'S dati Stlee1S, the AS4 OlX"e awears to be str<qer ttlan the T300 fibre no matter boW the strqI2l is measured, eittler the imp"e­gnated wor the sqte 0Iamalts ( see Table). 1M the atb!Jq ct the strqI2l values is 8Iso larger, ~ is 00IJtirmecS by the small value ct the Weibl11 m paramet« ct the strqtb <iistrirution det«mined from the ~ ct the strqtb w. gage lqUllineer r@tionsbipwitbio a ~ ct gdge Iqtbs . A eerioos involvtmeDt ct the small m pn­met« is the sIMp inaea ct the AS4 mean *. wit2l decr.q fibre ~ Jiqui­valent values ct m have been repmed in tbe receotlitsratlJre for AS4 /2-4/.

As stloWn in the Table, T300 anc1 AS4 am differ in diaIneter, SUlfa aspect and surface JRPDaticn The latter was iDitiat1y held to be very Jmpcltmt as carton fi1:res ordiDarely are stJWIIed sia9cS.

Me<t.anio:aJly ~ the 914 reem bebaw6 just1ik.e oUler tbennoeeCB currently used for stru::bJra1 <XlOlp06it:B8, more pWru1ar1y for britfieoess. But it has been verified several times that tbe medmlica1 p-q>emes ct tbe neatresin are fair1y depeodeot ()Q the thermal treatmentappBecS~the~cyde.A<XUlb!r4'Ject(11be fibre 00 the t1lermo-medlani p-opertieS (eg. the glass transi1Dl t:empefat1Jre (1 idlerIDa11y curec1 p-epreg <rupcrJS) was 8Iso ot&r ved . 'Illqb the same cure cyde was awlied to produre both the T300/914 and the ASlt/914 laminates, the identity ct the in situ matrix dlara<:t«isIia may be quesIiooab1e.

1.2 - IQta ()Q Jaminates

With respect to the fibres properties, the ~ ~ ( E. FJ was taken as the ratio (1 the oornpaiite rupture rocm to the fibre rupture &1Iain. strain beq not oonside­red to depend ()Q the actm1 fibre volume fraction.

695

As it oould be expected, r~ T300 by AS4 00es mainly affect the flb"e cootrclled ~<tthe laminates. Thisisquitevisitie\lor'hen the E. F.s <t the unidirec­tiooa1 oc the CfCES-}:iy 0-90-0 laminates are <:aI1pMed : it is obvious tbat the AS4/914 t:ensale Ft"opet1ies do not retle<:t the better Ft"opet1ies <t the AS4 fiNes. A Ft"oviiooal. cooclusioo. may be that there is a limitation to the ablity <t the 914 matrix to translate this better ~qth into the COOl}XlSit:e.

In <Xm~ the IX"~ <t the mUlt1dired1OOal 2~ laminates are moce r~U1 <t the fiNe ~ even if the 'bole ooetficient" ( notdled spedmenIpIain spedmen rupture strains ratio) is ~ in f~ d the T300/914 sy&em.

S:;me Ft"opet1ies are not considered to be <tiredly <mtrdled by the fittes : either they show no spedal tendency (eg. the int:edaminar shear strqth) oc they amnot be sim}:iy di9::usged due to the <mtributioo of irreIWcmtmecbanisms.lbatis the cas& foc the tooghne$ in mode I, measured 00. IXB specimeos W'iere the otmved R-rurves am 12ll1s Gl<: valUes are deeIiY inIlUenced by the un<m1rolled fmnation <t D«e ~ lSI.

A number d fads ~ cut ~ vifN1 : ~ a I:tdlBrili5e stress, in reference to the flb"e <tiredioo., is invdm, the AS4/914 S1*m pOOUces less good fJgllfes : transverse tensioo., tim;}:iy failure in the transverse layer of the Cf~ oc 28-}:iy laminates, fiDa11y the type <t rupture 00setwd in the OBrien edgt delanrinatioo. test. Thes& ~ticcs an tend to indicate tbat the p:x:t" quality of the AS/914 inbIdac» can be respoo.sible fcc the 1ack of transverse ~qIb.

1.3- Itrectmeasurement<t the interface IX"~

The interface IX"opet1ies have been detemIined ~ the micr<medmlical ~ of ~ carton fiNes embedded in the 914 resin by pun~ oc mvrlAGtation a<:XX:C~ to the IX"ocedures ~ have been deeaibed JftViouslY /fr 7/.

The pun-oot t£Ss p a similar, very hWl, ultimate shear ~qtb foc boUl fitces ~ttl coosiderable oca~ But the fragmentation t£Ss show tbat the shear 1Jan!ter ~ f1tce am mamiSbetter WlUlAS4due tottlelr bigber Wqttl atmt~. ~ the fl~ b!sts have been CIlly p!I1iIIy perf<llDed with the 914 resin jffitIlOW, tests with a IXJImA resin give q\Dt rW9antinf<l'!Datioosas far as fitce pq>er­ties are to be <:aI1pMed n I. NeiUler foc T300noc foc AS4, ~ isot.rged to oo::ur ~fragmentation ~<Xrlfirmsthe elDIIent ~ level am ~ aOideDo& to the pun-oot~qths.

FrOOl the IX"esent inwst¥.ticcs, it tmns out tbat the T3001 am AS't/914 inter­face are boUl <t ~ quality in sbear rut tbat AS4 may mwort a bigber md To give a fur1ller coofirmation, it is wa1b re<:aIliq!; tbat the mia~apJs d ~en T300/ am AS4/914 dry &pedmens dier quite similar Views with f~ cc no UD<X:IVered fitces ~

696

to the JX1ferentiSl <Xtlestvt rupblre in the ma1ltt a 1a)'W (( 1i¥tl1dl. ~ to itS granular appeaAIlCt, stm tldsts <Il fitns.

II - DI SCUSSION

~ the dirt<:tmeesurementot the interfada1 &bear strqUl (by p1ll-¢Ut) an:1 the &bear stress at m8ldmum of the ked transfer ( by tragmantab), a tentaUve inteqlretlt1on((theT300/anc1AS4J914medlanical~maybepltforward.

First, there is ant d8:X:rda1xe betw8en the * Iew1 d "adbesion-d both mres and therelative1ypo<:rstrqthd the AS4/9141amtDates in the stuatioIls d trans, .. Ioadq. NewrUleAees, one bas to be careful w1Ul the cooduIiOIJs of the medlanica1 t:eets whert a transverse tension is putoo tht plies : if tbe other materials tI6t8d in the present program art ampared, itis fooOO t2lat ~ Ume8outof tint, AS41amimttsar. bettIr in transverw strqtb (wlUl5208 an:15245 resms).

As for thtear1y format1ood the transvergeawin~ 1am1mte8, it bas been known for a kq time tbat the prooess stnqty depends 00 the ply t1'IDnee8 and the properties d tbe actjunt 1ayeB. Tatq tbe reie¥ant values for both matIria1s pes the usual trend d aadcs ~ mort and more ear1y as tht ~ ply ~ increase8 but the __ aa::t sensiIivit¥ d the AS4/914 matIria1 is also oonfir'Il*'.

Again, no man of this ~ sensitivity bas been foooo in the intQrface micro­mecbani<:al tBsts. yet. as reportsd by Nid):)IS 18/, teDsDt ~ in Ult transvenIt <Irec­Uon 'ftWId iOOuot, 12lroogtl the PasecIl ettect. a shear stress akq the fibre dnCticrl anc1 Weinberg, 00 glass fibre remtorced ~, pcints out t2lat the teDsDt strqt2l d the unidirectiooal transverse spedmen is in inVeBe prcpcrtion with the mean fibre Iqt2l d tht fully fragmented .. filament spedm8IlS 19/. Anyway, sudl a oorrtiation isd no ~ bert as the fibres art differtotin tmion.

In remrn, fcaowq a 61~ by Uematsu & aI./IOI, the merea d ., &bear stress at the m8ldmum of the ked 1ranster, 'Ift!len passmg fran T300 to AS4, can be dtrec­tty inv<lted to acmmt tor tbe poor K F. of tbe AS4/914 <XlmpC8t.es in t:e!lskXl : as it bas been otarved many urnes 00 m:xIel ~ \\o!lerl the fitlrelmatr1X ~ is very ~ anc1 the resin tlr1tUe, "penny Shape am-~ from the ems (( the tlr<*.erl fitns prq:egate ~ the reem 1bese aadcs are rwd8y viIibie in tht flagmented spedmens anc1 ~ modify tht IcedirC mode. The .. tht eoqy released by the titre rupture, Ult .. the aa::t diamet«, wbidl is just Ult case of Ult more reS­tantAS4 flbrt. In tht (! tensile loaded unidirediooa1 matsria1 an:1 rtlatld geometries, tht peony shape aacks occur UJl'Wlt to eedl brobn fibrt an:1 ooosequeoUy affect tbe doee ~fibres.

697

1be n&materla1 to break. in temi<Il ".,;n thm be the ooe with the moce resistant filns (AS4) due to a fa&1er f<mBticn d a weak sectioo. in the sb"essed specimens. Pm­thee mecDmism ~ be eo called upcm : as the aitic8111qth is ~ f« AS4, the ~fUesareover~overagreater1lqth. \Ir'bicb inaeeses the prob!ti­lily f« 'llfiM fibres to break. That be<xmes m«e prooouna!d with AS4 fibres of grGater variatlIity.

In!tS8d of the ~ suspected quality of the AS4/914 inteJfaoo, it is the oooperatiVe ~ of the rues and the resin \oIt'tD is f¥M in quesUoo : a very good fibre/matriX tx:IlcIq ~ with a lIP fibre sttqtb and Iiqv gcqe 1Iqth sensiti­Vity resuttin less "efficient" <XXDpOEIib!s as far as the lqitndjnal prq>erties are <XOa!I1led and the neJI:. tooome models of ccmpoaite rupture ~ to take it into 00DSiderati00. On the other band, the 1JaD&veJge ptpJtiI6 ~ bIftle the mia<me<ilarli<a inter­fac» ~ at the Dm8Dllt mu&t be <XDDied ttat parameWs different frem the ~ sbear sIrqth tested up touowgovem tbege prqlefties.

REFERENCES

1. GeJmm-{ren(h reeeerdl ~ Fioal Report (1966). 2. 1-M. Whitney & L. T.Ina1in "T~C<mposites",ASl'M SI'P 937,N. 1- joboston,

1icS7 ASIM,PbiIadeIJDa (1967) 179. 3- A T.aBenedetto, L.Nmais, L.~ & 1-Groeger, lstInbm Coot. 00. COOlp.

Interfaa!s am-i), H Isbm & 1- L. ~il.1!I!iIImec ~ Puit (1966) 47 4. P. W.Manders & T.W. Oloo. j. of ~PIasUcs & COOlp. 2 (1963) 43-5. D.Guedra.D. Lq.j.Roucml,C.MaraIs &P. SfIB4¥, ~ Jntcoo.t.oo.COOlp.Mat

Oa:M6.11md2), F. L. MiItthe\otr's & at.. il,1!Igerier Awl ~ (1967) vol. 3, p. 346.

6. M. S'anchez, G. Desarmot, M. .<:. Meriemle & B. Bamer, 5EImes joum8es Nat Composites ()t«:-5), c. Batbias & D. MenUs, il. &ditb1s PluraJis, Paris 0(66) p. 471.

7. 1-~.Fawe &D. ~ 6ttl Jntcoo.t.oo.COOlp.Mat.,q>.dt, vol. 5,p. 471. 8. D. 1-ttids in T.cmposite MatsriaJs : Testq: & Design", ASl'M SI'P 893, 1-M Whitney,

il. ASIM, PhfIadeJplIa (1966) 109-9. M. Wembe!gin ,.~ C<mpcBtes",ASIM SI'P 937, N. j. jcDlstal.1icS. ASIM,

~(1967)166. 10. T. UemaJsu, AMLmIuki, T.1keda & H Pumia in 'Ompa!ites'66:Reoent A<lvanres in

Japm & the Unbd Slates, Proc. Japm-U. S. aJ.f-III, K.. Ka\r.1ara & at..1icS. Tokyo 0(66)719.

698 BOO AS4

surface aspect striated smooth oxygen 1. (ESCA) 10 8 sizing type 5 no diameter (mm) 6.9 7.3

l2mm long single fibres Or (GPa) 3.83 4.45

€r (1.) 1.65 2.01

impregnated tows Or (GPa) 3.38 4.02

€r (\) 1.45 1.72 strength vs. length m 11.5 4.1

pUll-out 'Cu (MPa) 110-130 110-130 fragmentation (DGEBA) Ie (mm) 0.37 058

tm (MPa) 40.6 50.2

°rOe) (GPa) 4.30 8.03 fragmentation (914) Ie (mm) 0.33

'Cm (MPa) 53.3

unidirectionalO' OR (MPa) 1536 1545

€R (I) 1.09 1.09 E.F.( 1.) ref= tow 75.2 63.4 E.F.(\) ref= fibres 66.1 54.2

unidirectional 90' OR (MPa) 73 57

€R (\) 0.82 0.58

0-90-0 cross-ply lam. €R (\) 1.21 1.34 E.F.(\) ref= fibres 73.3 66.6

02-(90k02 cross-ply lam. initial state no cracks a few cracks 28-ply laminate .) plain OR (MPa) 625 826

notched (MPa) 437 552

€FPF (\) 0.64 052

EI .• efficiency factor (see text) ,,} quasi-isotrope (12-6-6-4)

Comparison of BOO and AS4 fibres and the corresponding laminates with Ciba 914

ENDOMMAGEMENT EN COMPRESSION ET EN TRACTION AUTOUR D'UN TROU D'UN MATERIAU COMPOSITE CARBONE/EPOXY

D. LAI, c. BATH lAS·

Universite de T echnologie de Compiegne Groupe Mecanique de I'U.A. 849 du CNRS - BP 649 - 60206 COMPIEGNE - France

·CNAM - 292 Rue Sf Martin - PARIS - France

ABSTRACf

The damage behavior of composite material under compression load is, specially when there exists stress concentration, a very important factor in its design. The present paper is aimed at the comprehension of mechanism and the characterization of failure in a carbon/epoxy composite plate. In order to the usual difficulties in compression investi­gation, the mechanical tests were performed with compact specimen. The damage evolu­tion was recorded to the variation of mechanical behavior, which yielded the criterion and law of mechanical behavior. Since the compression damage has a 3 dimensions form, the damage criterion were established in term of the energy by unit volume. The applications of these criterion shown that the linear fracture mechanic is a efficient tool for the charac­terization of the compression failure of the carbon/epoxy laminate materials.

INTRODUCTION

Les composites it fibre de carbone sont connus pour leurs proprietes mecaniques tres eleves. Mais ils ont aussi deux faiblesses patentes: la resistance a la compression et la resistance aux concentrations de contrainte, ce sont deux points qui ralentissent malgre tout Ie developpement de leurs applications dans I'industrie aeronautique. C'est pour cette raison que Ie comportement a l'endommagement en compression est un facteur tres im­portant dans la conception des materiaux composites, notamment lorsque ces derniers sont sujet a des concentrations de contrainte. Le travail presente dans ce papier consiste a etudier Ie mecanisme des endommagements agissant sur un materiau composite stratifie carbone/epoxy sous une sollicitation essentiellement en compression, et a caracteriser Ie comportement a ce type d'endommagements du materiau.

I - MISE AU POINT DES ESSAIS MECANIQUES EN COMPRESSION

Des essais mecaniques en compression ont ete realises sur des eprouvettes com­pactes specialement adaptees (Fig. 1). L'utilisation de cette nouvelle technique avait en fait pour but d'eviter des difficultes habituelles de I'etude en compression qui sont liees

700

aux problemes de flexion et de flambage. Pour eviter des endommagements inattendus, les charges sont appliquees par deux axes en acier enrobee de deux manches en acetal. Ces manches en acetal permettent de mieux repartir les efforts de contact et donc d'eviter un gradient de contraintes trop fort. Vne entaille en Vest amen agee It l'arriere de l'eprouvette de fat;:on It pouvoir placer un capteur de compliance et it eviter les contacts.

Au cours des essais mecaniques, l'ouverture de l'entaille de l'eprouvette compacte sont systematiquement enregistree, avec la charge appliquee, par un capteur It lame. Pour des essais de fatigue, les mesures sont effectuees periodiquement par arrets de machine et par une charge monotone. Etant donne que Ie materiaux etudie a un comportement tres elastique, Ie capteur d'ouverture n'est pas suffisamment sensible pour permettre Ie suivi de l'endommagement. Pour cette raison, un extensometre lateral avec un capteur it induc­tion magnetique, permettant de mesurer la variation d'epaisseur dans la direction de z, est place au bord du trou en sachant au prealable que l'endommagement s'amorcera It cet en­droit. D'autre part, une jauge de deformation est coHee verticalement sur l'axe symetrique de l'eprouvette. Afin d'eviter la saturation due it la concentration de contraintes creee par Ie trou, cette jauge est situee It une distance de 8,5mm du bord de ce dernier. Ces me­thodes de mesures sont realisables, grace it la nouvelle conception de l'essai de compres­sion: les eprouvettes compactes de compression ne necessitent aucun renfort mecanique, facilitent ainsi l'acces It ces dispositifs de mesures.

Les courbes enregistrees par les capteurs et la jauge sont comparees et correlees aux enregistrements des emissions acoustiques et aux resultats d'observations obtenus par rayons X et par coupes micrographiques.

Les champs de contraintes et de deplacements sont etudies analytiquement puis numeriquement au sens d'un etat de contrainte plane generalise, par la methode des "elements de frontieres" (These de LaY). Le milieu anisotrope et Ie rayon de courbure en fond d'entaille non nul sont pris en consideration dans les calculs. Pour faciliter les ana­lyses et les calculs, l'entaille de l'eprouvette compacte de compression, qui a une forme relativement complexe, est assimilee par une ellipse (fig. 1). La transformee conforme de variables complexes utilisee pour la formation de la fonction d'Airy est la suivante:

m· J a + ibf!i a - ibf!j' (1)

Dans Ie nouveau plan ~., l'entaille devient un cerc1e unite. A l'issu de cette analyse, et en considerant que l'endoihmagement se developpe de la fat;:on comme si c'etait Ie sommet de l'entaille qui avant;:ait regulierement (fig. 2), Ie taux de restitution d'energie pour une entaille et une charge donnees peut etre determine par l'integrale ci-dessous:

unite: J/m2 (2)

Etant don nee qu'en compression, la zone d'endommagement dans une piece en compo­site stratifie ne peut etre decrite dans un plan bidimensionnel, comme s'est Ie cas d'une fissure de traction dans une piece metallique, la restitution d'energie doit etre evaluee au terme de l'energie par unite de volume. L'integration s'effectue alors de A it B*. B* se situe sur la largeur fictive de l'endommagement, dj, qui est determinee par la repartition de contraintes a y dans Ie ligament de l'eprouvette et par la resistance it la compression du materiau, ayc (fig. 2). L'energie aura la forme:

unite: J/m3 (3)

701

II- RESULTATS EXPERIMENTAUX

2.1 Resultats mecaniques

Des essais de compression monotone et cyclique sont effectues sur les eprouvettes compactes. Pour verifier l'efficacite et la validite de l'utilisation de ces eprouvettes, cer­taines essais de compression sont egalement realises sur des eprouvettes rectangulaires donc plus "classique". D'autre part, afin de mieux comprendre Ie comportement it l'endommagement en compression et Ie mecanisme d'endommagement agissant dans Ie materiau, des essais de traction sont effectues, mais sur des eprouvettes CCT unique­ment.

Les resultats mecaniques sont resumes dans la figure 3.

La figure 4 montre des courbes typiques des essais statiques. D'une maniere gene­rale, en chargement monotone, aussi bien en traction qu'en compression, le materiau se comporte globalement elastiquement. Les courbes des essais statiques enregistrees par Ie capteur de compliance ou par Ie capteur de l'allongement ne presentent pratiquement pas de changement de pente jusqu'it la "rupture", bien que les premiers endommagements aient lieu bien avant. Autrement dire, ces deux dispositifs de mesure macroscopique ne sont pas sensibles aux developpement de ces endommagements statiques. Le capteur lateral et la jauge vertic ale places sur les eprouvettes compactes s'averent d'etre des moyens tres efficaces pour suivre l'evolution de l'endommagement. Sur des courbes correspondantes (fig. 4), des sauts en deplacement ou en deformation peuvent etre constates. Les cliches de radiographie X pris immediatement apres ces sauts et les courbes d'emission acoustique indiquent que ces derniers correspondent bien it la forma­tion des endommagements de compression.

En ce qui concerne les essais cycliques, des conclusions semblables peuvent etre obtenues (fig. 5). Macroscopiquement, la compliance des eprouvettes de compression n'est pas affectee durant les premieres 4/5 et 3/4 de la duree de vie respectivement pour un chargement Compression-Compression (C-C) et pour un chargement Traction­Compression (T -C). Cependant, Ie capteur lateral, la jauge de deformation et des cliches de radiographie X revelent Ie developpement et la croissance des endommagements des la premiere 1/5 de la duree de vie. Quant aux essais de fatigue traction-traction, des faits specifiques meritent d'etre soulignes. Lorsque les charges appliquees ont une amplitudes constantes, Ie developpement des endommagements autour du trou ou de l'entaille ne peut pas aboutir it la rupture finale ni a la degradation de la rigidite glob ale de l'eprouvette. Si les sollicitations component une amplitude croissante, l'eprouvette finit par se ramollir completement avec la generalisation de l'endommagement dans tout Ie volume. Mais la charge maximale peut atteindre 130% de la charge a la rupture statique de l'eprouvette! Ce phenomene, cons tate et confirme par de nombreux chercheurs est dG a I'effacement de l'effet de l'entaille par Ie developpement de l'endommagement autour de celle-ci.

2.2 Observations microscopiques

Des coupes micrographiques dans des plans successifs ont ete realisees sur des eprouvettes a differents niveaux de chargement ou a differents nombres de cycles de sollicitations. Le choix des plans de coupes a ete fait suivant des cliches de radiographie X de fa,<on a reveler significativement la forme de la zone endommagee. Les deux plans les plus interessants sont celui du ligament (A-A), et celui passant par Ie bord de l'entaille et qui lui est perpendiculaire (B-B).

702

Les resultats d'observations sont resumes dans Ie tableau ci-dessous.

COMPRESSION - 6crasement brutal de la zone avoisinant l'entaille MONoroNE - repartition irreguliere de delaminages dans l'epaisseur

- pas de rnicrofis lre COMPRESSION - pas de rnicrotlssure CYCUQUE - repartition parabolique de delaminages dans l'epaisseur REPETEE (C-C) - ecrasement des plis au centre de l'eprouvette COMPRESSION - rnicrofissures dans des plans de coupe /I au chargement CYCUQUE - repartition nes reguliere des delaminages REPETEE (T-C) - pas d'ecrasement TRACTION - endommagements antisymetriques CYCLIQUE - rnicrofissures dans toutes les directions REPETEE (T-T) - d6larninages dans les interfaces exterieurs

Les modes d'endommagement dependent fortement au chargement subis par les eprouvettes. Ces observations permettent d'effectuer des expertises sur des pieces rom­pues en service afin de remonter aux causes de l'accident. Avec les coupes microgra­phiques effectuees sur des eprouvettes avant et apres la ruine finale, nous pouvons conclure qu'en compression monotone, l'endommagement est essentiellement sous forme d'ecrasement de la zone avoisinant l'entaille, et qui a lieu brutalement au demier moment dans toute l'epaisseur. Cette conclusion permet de confirmer, en partie, la validite de l'hypothese prise sur la fa~on de developpement de l'endommagement, lors du calcul du taux de restitution d'energie.

III- CARACTERISATION DE LA RUINE DES EPROUVETTES ET CONCLUSIONS

L'evolution de la charge critique des eprouvettes de compression monotone depend non seulement du rapport a*IW, mais apparemment aussi du rayon de courbure en fond d'entaille p. Puisque pour une valeur de a*1W identique mais des valeurs de p diffe­rentes, les resultats sont disperses. D'ou la necessite de calcul de l'energie d'endommagement par unite de volume. Les resultats de calcul sont reunis dans la figure 6, ils sont tres satisfaisants. L'energie d'endommagement en compression monotone pour Ie materiau etudie est d'environ l,2·1Q7J/m3.

La figure 7 donne les courbes d'endurance du materiau sous sollicitations cycliques C-C et T-C respectivement. Le seuil de non endommagement (L*Jmax>Seuil pour un rap­port de chargement R=lO se situe autour de 5,3·1Q6J/m3, et celui pour un rapport R=-l est egale environ a 2,5·1Q6J/m3. Le seuil de non endommagement est nettement plus faible dans Ie cas de fatigue T-C que dans Ie cas de C-C. Le chargement cyclique en trac­tion-compression est donc Ie mode de sollicitation Ie plus penalisant pour les materiaux stratifies, du moins pour Ie materiau etudie. Compte tenu des mecanismes d'endommagement reveles plus haut, ce resultat semble logique. A souligner que dans les courbes de la fig. 7,l'energie L'imax, calculee avec la formule (3), correspondant au taux de restitution d'energie initial, n'a qu'une signification indicative. Car Ie mecanisme d'endommagement et la forme de la zone atteinte ne sont pas du tout les memes en solli­citations cycliques qu'enchargernent monotone. L'energie d'endommagement en fatigue doit etre calculee en 3 dimensions, par des integrales triples, en tenant compte de la geo­metrie complexe de la zone endommagee (These de Lai). Toute fois, il existe un rapport entre Ie seuil de fatigue et l'energie de 'rupture' statique, il est donc possible d'evaluer la duree de vie avec Ie parametre L*Jmax calculee avec la formule (3).

703

..

"'"

Fig. 1 - Eprouvette compacte de compres­sion et dispositifs de mesures

Fig. 2 - Hypothese sur la progression de l'endommagement

1,0 .---t""'"""_---------------, __ .....,..,....__---------,

0,5 compression

E

~ 0,0 ~-"""""-__ ....:... t:>

-D,S monotone

-1,0 L-_____ ==-________ --=--.:..::...:..:.... _____ --..l

Fig. 3 - Resume des resultats d'essais mecaniques

~-------------~O

Compression direcle Ep. compacle a*=34 p=4

L-'---'----'--"'----'-----''---....L..---'--"'-----'----'O -1,5

-2 Ouvenure (mm)

Fig. 4 - Essais de compression monotone

+0,1 Variation d '~paisscur (mm) 0

C jauge

Compression par ~tapcs Ep. compacte a*=34 p=4

z 0 ~ .., "" a

.z:: U

L-------------------~ - 15 o ' Deformation Jauge (x J 0'))

-2,4

704

o e-/p-~ 8-/p-

a CICf '. ,0 (,j."j

, U U

0,8 ,~

~ E

0,6 5 c:

" u

0,4 .~ 'E. §

0,2 u

0,0 0,0 012

Nombre de cycles (x105)

o~=-~~==~~~ ° 1 2 3 4 Nombre de cycles (xl05 )

Fig. 5 - Courbes d'essais cyc1iques Compression-Compression et Traction-Compression

1,6

1,4 • ----~--~--~----------~------~1,2 .....

1. I ° u ... 0

..... ' '30,8

>< p''-;:0,6 .-...l0,4 o Ep. compacte - L'l~moy= 1,138

• Ep. rectangulaire - - - . .c

0,2 L lcmoy= 1,257

12 11 10

.::'9 & 8 ~ 7 ~6 II 5

*J~ 2

\ \

R=1O 0 Ep. compacte • Ep. rectangulaire

R=-I • Ep. compacte . ~

~-----~------o~~~~--~~--~--~~ 0,0

0,0 0,1 0,2 0,3 0,4 0,5 0,6 104 loS 106 107 108

a*/W Nf

Fig. 6 - Energie de "rupture" en compression Fig. 7 - Endurance du materiau

REFERENCE

1 - Adsit N. R., Compression Testing of Homogeneous Materials and Compo­sites, ASTM STP 808, (1983) 175-186

2 - Bathias c., Esnault R. and Pellas J., Composite, July (1981) 195-200 3 - Lal D. et Bathias c., 5eme Journees Nationnalles Pour les Composites, (1986)

179-190 4 - LaY D. and Bathias C., Fifth International Conference for Mechanical Behaviour

of Materials, Vol. 2, (1987) 1231-1238 5 - Lal D., These, Universite de Technologie de Compiegne (1988) 6 - Lamothe R. M. and Nunes J., Compresion Testing of Homogeneous Materials

and Composites, ASTM STP 808, (1983) 241-253

221 7 - Matondang T. H. and Schlitz D., Composite, Vo1.15, No3, July, (1984) 217-

8 - Tsangarakis N., Journal of Composite Materials, Vo1.19 January, (1984) 47-57 9 - Wang A. S. D., Composite Review, Volw. 6, No 2, Summer, (1984) 45-62

THE INFLUENCE OF SPECIMEN GEOMETRY AND TEST CONDITIONS ON THE TENSILE

AND FRACTURE MECHANICS PROPERTIES OF GRP

A.V. LIMA, J.A.O. SIMOES', A.J.M. FERREIRA', A.T. MARQUES', P.M.S.T. de CASTRO'

ABSTRACI'

Universidade de Tras-os-Montes e Affo-Douro Av. Almeida Lucena 1 - 5000 VILA REAL - Portugal *Faculdade de Engenharia - Universidade do Porto

Rua dos Bragas - 4099 PORTO - Portugal

In order to develop data and tools useful for the design of structures made of GRP, a programme of mechanical testing and a finite element code are currently being developed. An outline of the work conducted so far is presented.

The influence of size, specimen geometry and applied strain rate on the mechanical properties of GRP has been studied. The fracture toughness of GRP has also been determined. Kc has been determined using SENT (single edge notched tensile) specimens and plates with a central notch. Three different types of K calibrations were used for the notched plates.

The behaviour of this material in the tensile test, although non-linear, may be approximated by two straight lines with different inclinations. A simple finite element programme was developed to model this type of behaviour, and the study of a notched plate is given as an example of application.

INTRODUCTION

Studies carried out by Rebelo et al [1) and Ferreira et al [2) have shown a linear variation of rupture strength and modulus with the area of cross -section for GRP samples. It is apparent that above a given cross section the modulus and tensile strength are geometry independent.

Another important aspect in composite materials of polymeric matrix is the possible strain rate sensitivity of the materials, and this has also been studied for cross-head speeds varying between O. 2 and 900 mm/min.

706

The inherent deffects present in GRP give support to the use of Fracture Mechanics as a possible tool for design of structures made of composite materials. Hence the critical value of the stress intensity factor Kc was determined for GRP samples using two types of specimens (SENT and central­notched plate), with three different K calibrations for the plate specimens.

To model the mechanical behaviour of GRP laminates, where the reinforcement is CSM-chopped strand mat, a linear finite-element computer programme has been adapted for non-linearity using a simple two-straight lines approximation to the stress/strain curve.

I - EXPERIMENTAL PROCEDURE

1.1 - Mechanical Properties

The studies were carried out with an isophtalic polyester resin (CRYSTIC 272) reinforced with chopped strand mat (600 g/m2) having an average weight fraction of 37%. The samples were cured during 24 h at 21°C, and had a post­cure of 3 h at 80°C.

The specimens used for the tensile tests were parallel bars 250 mm long, width varying between 10 and 40 mm, and nominal thicknesses of 3, 4 and 6 mm.These tests were carried out with an lNSTRON 1125, at 21°C and cross-head speed of 2 mm/min. The results are presented in figs I and 2, which also include a comparison with results obtained by Ferreira et al [2].

It has been concluded that for thicknesses over 6 mm, the mechanical properties are geometry independent. However, for thicknesses less than 6 mm, a minimum cross-section area of 100 to 140 mm2 is required to have geometry independent properties.

As far as the influence of cross-head speed is concerned, the results obtained showed that the tensile strength is approximately independent of cross-head speed, whilst the modulus and elongation at rupture are strongly dependent above 100 mm/min, fig.3.

1.2 - Fracture Mechanics

The same type of material was used for these studies, with an average weight fraction of 40%.

SENT specimens, with crack length/width ratio 0.45 < c/W < 0.5, were tested. The stress intensity factor calibration used was taken from Brown and Srawley,[4]

KI= (P c l12 IBW)[1.99- 0.41 (c/w) +18.70(c/w)2_38.48(c/w)3 + 53.85 (c/w)4] (1)

where P is the maximum ioad, Band Ware the thickness and width of the specimen, c being the crack length.

For the plate with a central notch three different calibrations were used:

KI =( P/BW) ( 7t c)1!2, solution for the infinite plate,[4] (2)

707

K 1= (P c 112 I BW)[ 1.77 + 0,227 (2c/w) - 0.510 (2 c/w)2 + 2.700(2c/w)3] (3) taken from Brown and Srawley,[4]

KI = (P/BW) [It (c + ao ) ] 112 with ao = 4 mm obtained by Daniel [5], (4)

In equations (2) to (4) c is half-crack length and ao (last equation) is the damage zone dimension.

Fig. 4, plotting of Kc as a function of B for SENT specimens, shows the independence of Kc. Slightly smaller values of Kc were obtained for lower thicknesses but that is attributed to lower fibre weight fraction.

In fig. 5 the results obtained with a plate having a central notch are presented and the calibration given by equation 4 shows an independence of mode I fracture toughness Kc with respect to 2c/w.

2 - FINITE- ELEMENT MODELLING

A finite-element programme described by Fenner [6], which uses CST -constant strain triangle elements for bi-dimensional analysis was modified to simulate non linear behaviour, and applied to the study of a GRP plate (quasi-isotropic) with a central notch (hole). The stress/strain curve (fig. 6) was approximated by two straight lines. This simple simulation of non linear behaviour is based upon the following procedure: (i), the stress state is computed for the load of interest P, using EI, and the Tresca equivalent stress

cr eq in each element is determined. For all the i elements where cr eq > cr m ( cr m corresponding to the intersection of the two straight lines), the following calculation is carried out

(5)

P. values are arranged in an ascending sequence, the lower value b~ing P ; (ii), the programme is run with P. = PI using EI since for this calculation all elements have a ~ a : Then, (iii), an iterative procedure is carried out, where for p~ I tWe programme is run for ~P = P.+I - P .. The computed stresses 1 Xnd strains are added up to the previoijs valu~s, and for each element where a is exceeded, E is used The iteration ends when the summation of load~ equals the ori~inal va­lue of P.

Fig. 7 shows a detail of the mesh near the zone of interest, where it is possible to appreciate the shape of the damage lone. In this example, the maximum equivalent stress is near the ultimate tensile strength of the material The finite element code developed is suitable to the assessment of the behaviour of plane structural details of interest, such as stress concentration situations.

3 - CONCLUSIONS

The following conclusions could be drawn so far:

708

i) For the material studied, in thicknesses lower than 6mm, mechanical properties are independent of specimen size if cross section is greater than approximately 120 mm2.

ii) Tensile modulus and elongation at rupture are strongly strain rate sensitive.

iii) The toughness Kc is thickness independent for SENT specimens.

iv) The calibration for a plate with infinite width, corrected for damage zone, gives Kc independent of 2 c/w.

v) The approximation of stress/strain curves by two straight lines and the use of a simple non linear aproximation describes with a reasonable accuracy the behaviour of a GRP plate with a central hole under tensile load.

REFERENCES

[1] - Rebelo, C. A. C. C., Ferreira, A. J. M., Marques, A. T., de Castro, P. M. S. T., "The Influence of Processing Conditions on Mechanical and Fracture Properties of GRP Plates", in:" Mechanical Behaviour of Composites and Laminates" ,W.A.Green, M.Micunovic,eds.,Elsevier Applied Science, 1987, pp. 54-63.

[2] - Ferreira, A. J. M.,Marques, A. T., de Castro, P. M. S. T., Rebelo, C. A. C. C., " A Influencia da Geometria dos Provetes nas Propriedades Mecinicas de PRFV", Materiais 87, SPM, Braga 1987.

[3] - Maciejczyk, I.A. , Slobodzinski, A.E. , "Guidelines for the Generation and Use of Data" in Military Handbook 17 A Part I. Plastics Technical Evaluation Center. January 1977.

[4] - Brown, W.F. , Srawley, J.E. , "Plane Strain Crack Toughness Testing of High Strength Metallic Materials", ASTM STP 410, 1966.

[5] - Daniel, I.M., "Mixed-Mode Failure of Composites Laminates With Cracks". Experimental Mechanics, December 1985, pp. 413-420.

[6] - Fenner, R. T. , "Finite Element Methods for Engineers" Macmillan, 1975 . • Ref. 2

erR (MPa)

100

· ",m '""'~i',~m~ ~.,;"

-'

Fracture Energy (K J,.,;2)

60

. 50

" , , • ~/.O ~ 1 1 JO

20

500~----2~0----4~O----~60----~80~---10LO----,L20~10

AreA (mm') Fig. - 1 - Tensile strength and energy aL rup-cure/res'1stant area

for a typical CSM polyester laminate.

'" w

ro IL

"' ~ 0

0

~ :L

Modulus (GPa)

10

9

8

7

6

5

4

_e--

•

• R~f. (2) o nom.lhickn.

3.8mm

-+1 __ --r---r ----.!-'---;- - •

4

oL-____ ~ ____ ~ ______ ~ ____ ~ ______ L_ ____ _L~O

5

3

o 20 40 60 80 100 120 Ar~a (mm')

Fig. L - Young's modulus and strain at rupture versus resistant area for a typycal CSM polyester laminate.

Modulus

ER

log " (mm/min)

Fig. 3 - Strain rate dependence of Young's modulus and elongation at rupture for a typical CSM poly­ester laminate.

709

710

~

N

~ 'E .. 0-1:

u ::s:

14

N 12 ..... 'E "

10 0-X

8 0 ~

6

4

15

10

0

2 4 8 12 16 Thickness (mm)

?ig. 4 - Kc ~:!~~:t~~ickness for a typical CSM polyester

•

0.1 0.2

2c/W 0.3

0--- eq.(2) .-- eq. (3) .---eq.(4)

0.4

Fig. 5 - K versus 2c/w for a typical CSM polyester c laminate_

0.5

N

'E :z ::[

100

~ am I/O

~ 50 ... VI

o 0.5 1.0

Strain ('/,)

1.5

E1= 8450 MPa

E2 ~ 5000 ~1ra

C1m!:: 60 MPa

V = 0 . 26

11g. 6 - Tensile stress-strain curve for a typical CSM poly­ester laminate approximated by two straight lines.

damag~

,

-J-.-

Fig. 7 - Detail of the mesh used for the analysis of a notched plate, displaying the damage zone.

711

SEMI·EMPIRICAL MODELLING OF STRESS RUPTURE DATA ON GLASS REINFORCED PLASTICS

J. SILLWOOD, J. AVESTON

National Physical Laboratory - TW11 OLW TEDDINGTON MIDDLESEX - England

ABSTRACT

For cost effective design of GRP composite structures in aggressive environments it is important that realistic safety factors are used which are based on testing experience and an understanding of the mechanisms affecting long term failure. Accepting the complexities of defining an accurate model suitable for all GRP systems, this paper utilises an expression for failure time of stress rupture specimens based on slow crack growth, which has been experimentally verified for two extremes of composite behaviour. The approximation is shown to give an acceptable fit to unidirectional and CSM glass composites in water at ambient temperature.

INTRODUCTION

To enable the designer of engineering structures to determine safe and efficient working loads for glass composites that may be exposed to aggressi ve environments, it is desirable that a characterisation of the failure process for extended lifetime measurements should be obtained as a basis for extrapolation.

In order to utilise the strength of glass fibres in GRP, fibre lengths are chosen so that failure usually only ensues when fibre fracture occurs. It is well known that glass fibres are weakened in environments of water and dilute mineral acids and it has been shown/if that glass fibre degradation is the dominant mechanism of weakening for GRP materials in these environments when the fibres are aligned in the direction of stress.

Models based on subcritical crack growth attributed to Charles/2/ have been used to describe the failure of two extremes of unidirectional composite failure/!,3-5/ and this paper attempts to generalise the expression for time to failure used in these models to

714

give an approximation for application to other composite systems between these two extremes.

1 - TIME DEPENDENT FAILURE

Recent attempts to model the long term failure process of glass and ceramic materials in aggressive environments have used the principle of sub-critical crack growth utilising the formulation based on stress intensity following Davidge et al/6/

ie:

II = AKn I

(1)

where II is crack veloci ty , K is crack opening (mode I) stress intensity and A and n are materials constants. It is assumed that the crack grows under the influence of the environment and applied stress according to equation 1 until such time as it reaches a critical size, at which point catastrophic failure occurs. Integration of this equation gives an expression for time to failure as a function of applied stress, the relevant characteristic constants for the material and testing geometry.

1.1 Glass fibre bundles

Using the above equation for crack growth, Kelly and McCartney/7/ have given an expression for time to failure of an uncoupled strand of glass fibres assuming that pre-existing defects are present in the glass fibres and that their individual strengths can be described by a two-parameter Weibull distribution. The resulting expression they derived for time to failure contained an integral which could only be solved numerically. It was shown that the resulting curves for time to failure at different values of the crack growth parameter naIl converged asymptotically to the short term strength of the bundle. They went on to show that at long times to failure when the applied load F was small the time to failure depended on the applied load to the power -n, the crack growth parameter in equation (1) ie:

t a 2-f ~

(2)

Their model was validated experimentally/1/ using strands of E-glass fibres loaded continuously in water until failure, although the crack growth rates in E-glass fibres were beyond the range of growth rates measured for bulk glass from which the constants A and n were calculated.

1.2 Monolithic crack growth in GRP

In extremely aggressive environments it has been observed that cracks propagate in GRP, possibly from fibres exposed on the surface, and pass into the bulk of the material uninterrupted by interfaces between matrix and fibres. The crack surfaces produced are planar, particularly at long times, and the crack grows until such time as

715

catastrophic failure occurs when the critical stress intensity is reached at the tip of the largest crack. By assuming a crack growth rate according to equation 1 and integrating, an expression for time to failure can be derived/3/, resulting in a similar expression to that previously used for time dependent fracture in bulk glasses and ceramics, ie:

2 [K 2-n _ K 2-n] 0 2 y2A(n-2) Ii IC

a

(3)

where a is applied stress, KI . is the initial stress intensity, Y is a geom~tric factor anf KICl. is the c1itical stress intensity. Substituting KIl.' a Ya. anQ~IC a Ya. in equation (3) gives, a l. max l.

2 a t f (1 - (_a_)n-2) (4)

an yfi A (n_2)a. n/ 2- 1 a max a l.

where a. is the initial crack size and a is the short term strength~ max

It was found/3/ that equation (4) fitted the results for stress rupture tests on unidirectional E-glass GRP in dilute sulphuric acid using constants derived from dcb crack growth measurements and assuming an initial crack size of - 0.14 mm, roughly comparable with the strand dimension.

Equation (4) has two important features that also appear in the analysis for uncoupled glass strands. At short times there is a limiting value of stress approaching the short term strength and at long times to failure where aa « amax ,

t a _1_ f n

a a

1.3 Generalised theory

(5)

The two regimes described in 1. 1 and 1.2 can be considered as representing two extremes of composite behaviour. In the first case there is no interaction between adjacent failures in the glass fibre bundle, ie no stress transfer, and in the second case the failure process is completely dominated by interaction between nearest neighbours. ie failure occurs by the growth of a single dominant crack.

The analyses describing the stress rupture behaviour of both of these regimes have common features, ie a limiting short term strength and a straight line log-log relationship between time to failure and applied stress at long times where the stress is small compared to the short term strength. It would therefore be reasonable to assume that systems lying between these two extremes would also follow a similar pattern. In support of this assumption, it is observed that the fracture surfaces of CSM in stress rupture experiments in water do

716

become progressively more brittle as the loads decrease and the failure times become longer, and of course at short times the strength must reach a limiting value.

To provide an approximate characterisation of this intermediate behaviour we return to equation (4). By combining some of the constants appearing in the equation which refer either to the material or the mode in which the specimens are loaded, a simplified expression for time to failure can be written, ie:

with B = 2

yfiA(n_2)a. n/ 2- 1 1

B a t = (1 - (_a_)n-2) (6) f n a a a max

which contains the essential features at short and long times to failure.

2 - EXPERIMENTAL

Equation (6) has been tested for application to a number of GRP systems loaded continuously in water, including:

i) unidirectional orthophthalic polyester impregnated E-glass strands with a nominal volume fraction of 50%.

ii) unidirectional orthophthalic polyester E-glass composites 1 mm thick with a nominal volume fraction of 50%.

iii) orthophthalic polyester E-glass CSM 5 mm thick with a nominal volume fraction of 20%.

Details of material manufacture, specimen geometry and testing procedure can be found in references 1, 3 and 8.

3 - RESULTS

Using a non-linear optimisation routine/9/ on an ICL 2972 computer the stress rupture results were fitted to equation (6) by selecting values of the parameters B, n and a such as to minimise the sum of the squares of the deviations, ie a~ast squares analysis.

The results for the three systems used to test the analysis are shown in Figures 1, 2 and 3. In Figure 1 the Y-axis is expressed in terms of applied load in kgf rather than stress as used in Figures 2 and 3. Although there is a degree of scatter in the results, particularly for the CSM at short times, the data fit reasonably well to equation (6).

In order of severity the impregnated strands fall in strength most rapidly, followed by the unidirectional GRP and the least

717

severely affected material is the CSM.

4 - DISCUSSION

The ability to be able to arrange the fibres in composite materials in such a way that the properties are optimised for a particular application leads to a diversity in structure which influences the response to mechanical loading, eg compare the tensile load/deflection curves for unidirectional and 0/90 laminates. In the case of the unidirectional material the curve is approximately linear to failure while the curve for the 0/90 laminate begins to deviate at a strain corresponding to the strain at failure of the weakest point in the 900 phase. Other forms of GRP such as CSM and woven glass materials will behave in a similar manner to the 0/90 laminate in that cracks will initially occur in areas where fibres are not aligned in the direction of loading. The spacing and separation of the surfaces of these cracks will be dependent upon the structure of the material, eg lamina thickness, strand size in CSM, level of coarseness of warp and weft in woven materials, and the magnitude of the strain. This damage is also likely to affect the speed at which a liquid environment can penetrate into the bulk of the material and hence affect the speed of any subsequent weakening that the environment may cause. For the above reasons, it follows that the task of defining an accurate model to describe the loss in strength with time under load in an aggressive environment for all categories of GRP would be extremely complex. Subtle changes in the model would need to be made to take into account the changes in the behaviour of each type of material, but the high degree of scatter in strength measurements of many forms of GRP would not facilitate reliable validation of any particular changes that might be made.

It has been shown that, although the model described here is only an approximation, the data have fitted reasonably well for two unidirectional materials and for a CSM material with different optimised values for the parameters n and B. The significance of each of these values in terms of physical meaning is not clear in that the parameters are not independent of each other during the optimisation, and serve only to describe the experimental data in a form that can be given credibility for two extreme composite configurations.

At long times to failure the model approximates to a linear log-log relationship, and it is tempting to suggest that this would suffice. But unless account is taken of the short term limiting strength it would not be clear which data points should be included in the fitting procedure as the applied stress increases and any errors would affect the fit at long times.

The procedure is only valid as a three parameter fit if the data points give a clear indication of a limiting strength and tests for significance have been written into the optimisation routine. If the data are all at relatively long time, errors in estimating the limiting strength will occur, particularly if failures at the higher loads occur in an unexpectedly short time. In this case a two

718

parameter fit for B and n would be a better solution if the short term strength can be measured independently. An example of this is shown in Fi~re 4 where data from a mixed woven roving-CSM material in water at 40 °c from Phillips/l0/ has been tested with both a three and a two parameter optimisation. The data are averages of between five and about twenty individually recorded times to failure and are all at relatively long times. The three parameter fit is sensitive to the single point on the extreme left and uses this to predict an underestimate of a ,in this case an ultimate tensile unit load (UTUL). Using the ~Wo_farameter fit and the independently measured UTUL value of 256 N mm this point becomes less significant.

It must be accepted that the present approximate model is deficient in a number of respects. It does not take into account the possible effects of diffusion of the water into the matrix which would appear to be significant if the two sets of data in Figures 1 and 2 are compared. The fall in strength for the smaller impregnated strands is more severe than the bulk unidirectional material, suggesting that there may be a size effect, allowing faster saturation in the impregnated strands and resulting in more rapid glass degradation. Some progress has been made in trying to account for the effects of diffusion/ll/ in unidirectional material assuming that the diffusion process is unaffected by the application of load, but for CSM material this will clearly not be the case.

The model also assumes an initial constant crack size which again is not the case for CSM. Roberts/5/ has used a similar analysis to interpret the results of constant load bend tests on CSM laminates in hydrochloric acid, but also took into account a stress dependent initial flaw size. However, the change in stress distribution in the bend test occurring as a result of cracking did not appear to be taken into account.

Finally the controlling growth law is assumed to be constant throughout the range of stress applied and evidence to suggest that this may not be assumed has been given by Phillips/l0/ at temperatures of 40 °c (see Figure 4) and 60 °C. He observed that at short times to failure the stress rupture curve was unaffected by temperature, but at longer times a change in shape was detected.

5 - CONCLUSIONS

It has been shown that a generalisation of an equation used previously to explain the results of GRP stress rupture experiments where growth of a single crack predominates can be used to describe the room temperature results of other systems where failure occurs primarily as a result of stress corrosion of the glass fibres. Further refinements to the model are possible but this would require detailed knowledge of the effects of diffusion on long term strength, the severity of microdamage with increasing stress and the effects of microdamage on diffusion. At higher temperatures, however, previous evidence suggests that a more sophisticated model may be required.

.., o .2 .., .!! 0-0.

15r------------------------------------,

00

Measured 0

Predicted --

« 5

lyr 5yrs

OL-______ ~------~------~~L-~

Ln (time to fail)

719

Fig 1 - Stress rupture in deionised water at ambient temperature of E-glass strands impregnated with orthophthalic polyester resin

1200,-----------------------,

1000

200

M .. asur .. d • Pn~'dicl.d-

10 15 Ln (tim. to fait)

1yr Syrs

20

Fig 2 - Stress rupture in deionised water at ambient temperature of 1 mm thick E-glass/orthophthalic polyester resin composites

720

120,----________ -,

1()()

'" 80 'E z ! . . ~ 60 ;;;

" .~ 0. a. .. 40

20

Fig 3 -

Measured •

Predlcted-

. . . . :

Iyr 5yrs

10 15 20 Ln (time to fall)

Stress rupture in deionised water at ambient temperature of 5 mm thick CSM E-glass/ orthophthalic polyester resin composites

REFERENCES

"§ ~ a.

" ~100 " .~ 0. a. ..

50

Measured •

Predicted-

Ln (timE" to fail)

Fig 4 - Stress rupture in distilled water at 40c of mixed E-glass woven roving/CSM/isophthalic polyester resin compo­sites from Philips/l0/

1 - J. Aves ton , A. Kelly and J.M. Sillwood, 3rd Int. Conference on Composite Materials (1980) 556-568, Paris.

2 - R.J. Charles, J. App. Phys. ~, (1958) 1554-1560. 3 - J. Aveston and J.M. Sillwood, J. Mater. Sci. 17, (1982)

3491-3498. 4 - D. Hull and J.N. Price, J. Mater. Sci. 18, (1983) 2798-2810. 5 - R.C. Roberts, J. Mater. Sci. 20, (1985)~341-1350. 6 - R.W. Davidge, J.R. McLaren and G.Tappin, J. Mater. Sci. §,

(1973) 1699-1705. 7 - A. Kelly and L.N. McCartney, Proc. R. Soc. Lond. A, ill,

(1981) 475-489. 8 - J. Aveston and J.M. Sillwood, NPL Report DMA(A)46 (1982). 9 - A. Woolf, NPL Report DMA(A)79 (1984).

10 - M.G. Phillips, Composites 14, (1983) 270-275. 11 - J. Aveston and J .M. Sillwood, Trans. 1. Mar. E. (C) 91,

Conf.2, (1984) 143-147.

ABSTRACT

A MODEL OF LAMINATED COMPOSITE PLATES ASSURING THE CONTINUITY OF DISPLACEMENTS

AND TRANSVERSE SHEAR STRESSES

M. TOURATIER', Q. LlUT, P. LORY"

'ENfT-Laboratoire de Genie de Production BP 1629 - 65016 TARBES - France "RNUR - Direction de fa Recherche Service 0861 - 92508 RUEfL - France

In this paper, we describe a two dimensional model of plates which takes into accountca parabolic variation of the transverse shear strains through thikness, so that there is no need to use shear correction coefficients in computing the shear stresses. These properties are very important for the modeling of moderate­ly thick laminated structures. The contribution of the present work is to take into account exactly the contact between the layers of the structure. The present method uses a classic dis­placement approach for higher-order shear deformation. The form of the displacement field is dictated by the satisfaction of the following conditions: i) continuity of displacements and stresses at the interfaces ; ii) disappearance of the transverse shear stresses on the plate surfaces while non-zero elsewhere. This re­quires the use of a displacement field in which the inplane dis­placements are expanded as cubic functions of the thickness coor­dinate and the transverse deflection is constant throughout the plate thickness. Consequently, the normal transverse deformation is discarded, because we want to have a displacement field which contains the same dependent unknowns as in the first-order shear deformationtheory. So, it is possible to obtain numerical results by a Cl finite element approximation for the displacements. In the same way, a higher-order model with the five classic inde­pendent generalized displacements, and which takes the normal transverse deformation into account, involves a C2 finite element approximation for the displacement /1/.

722

SUMMARY A great number of plate theories exist in the literature /2/.

First, for thin plates the Kirchhoff - Love theory in which the transverse shear deformation is discarded. An adequate theory for layered composite structures must account for transverse shear deformations in order to account for the shear stress transfer at the interfaces of the layers. In Hencky - Mindlin - Reissner -type theories, that is the first order shear deformation theory, the displacement field accounts for linear variations of midplane displacements through thickness. In addition, these shear defor­mation theories do not satisfy the conditions of zero transverse shear stresses on the top and bottom surfaces of the plate, and require a shear correction to the transverse shear stiffnesses. Some works are devoted to these shear correction factors, see for instance /3/. In another way, the higher-order sheardeforma­tion theories of Ambartsumian /4/, Kakowsky /5/, Panc /6/, Reissner /7/ do not account for displacement and stress continuity conditions at the interfaces between layers of the laminated struc­tures, while the three dimensional theories /8/9/ of laminates are intractable as the number of layers becomes moderately large. Fi­nally, a two-dim'ensional theory of plates that accurately descri­bes the global behavior and the stress distribution for laminated plates is very convenient in the use of finite element method in engineering applications. The work summarized in this paper con­cerns the modeling of moderately thick laminated structures by using both analytical and numerical approaches. The analytical approach allows us to build an accurate kinematic description for the laminated plates, while the finite element approximation allows ~s to analyze complex laminated structures by development of a finite element based on this kinematic description.

Consider a plateIl of uniform thickness h in coordinate system O/x,y,z; z being normal to the plate's mid-surface, and Min (Max IIxU , Max Uyll ) » Max (z). The plate is assumed to be subject to surface load q acting along the z-axis at z = h/2 and tangent

loads will not be taken into account on faces z = ~ h/2. Because h is small in comparison with the other dimensions of the plate, and since the normal transverse strain is neglected, we can write the kinematic as follows, for each layer I of the plate in the coordi­nate system O/x,y,z :

tU ... = lA.oI ~l. ~oI" ~2.' ~,,~ ) al:lC. or':i; t U!o = W (~) where u~are membrane displacements, w the deflection. The genera-lized displacements . u ... , w, LII)", I tWoel ,are functions of x, y and if necessary of time. From a complete three-order expansion to the power of z for an homogeneous plate, it is easy to verify that ~01 -= 0 .Wi th (1), we want to.satisfy the continuity conditions for disp'acements and stresses by a displacement approach. To summarize the method, the-description of the model is limited to the uncou­pled membrane-flexion configuration. So, the constitutive law fo~ shear is (for each layer 1):

723

.t t l to 1 0"' :=. G (to ... ; 1. w -'to W)rX. ) \2.)

et 0( "''I III~ , / ...... ,·CXE,'x,':}L without summation on cI. and with W,aC. = OW u..... ,. J

and where t~. is the shear modulus of the layer 1 for the stress t~ • The continuity conditions for displacemems and stresses

are respectively at the interface between the layers 1 and 1 + 1, of the altitude zl ~ 0 :

1+1 1+1 From (3) and (4), it is easy to eliminate c.u rio 1 and W ~3

It is sufficient to extend these continuity conditions (3) and (4) at all interfaces of the plate to obtain as unknowns only w)rJ.. ; ·w.~ ; ·w",~ , where I = 0 is the reference layer.

Finally, the unknowns ow",,, are eliminated by the satisfaction of the conditions of zero transverse shear stresses on the top and bottom surfaces of the plate, where we must have :

t",TO\> ) I ("P..ott.", ) (1'1) 0 ( cr:. l."d:h/:. = \ e:.. "2. ... :th/oz. = ~ l.::Hlh."=

(5)

SOME REMARKS The advantage of this model is that it contains the same dependent unknowns as in the first-order shear deformation thoery, but ac­counts for parabolic distribution· o.f the transverse shear strains through the thickness of the plate and the continuity conditions for displacements and stresses through the interfaces Of the non­homogeneous structure without shear correction factors. This model allows us "to relocalize" easily for stresses and dis­placements in each layer from (1) and (2) satisfying (3), (4), (S). To apply this model to the analysis of the complex composite struc­tures, a C1 triangular finite element has been developed /10/. The element has twenty sevendegrees-of-freedom located at the corner and mid-side of the triangle. Several numerical tests have been carried out, and have given good results for both thin and moderately thick plates on the standard tests and in comparison with the exact tridimensional theory for laminated orthotropic plates /11/.

ACKNOWLEDGMENT

This work is financed by Renault, Direction de la Recherche, under the contract Nr SERAM/30.

724

REFERENCES 1. M. Touratier, Mechanics Research Communications, 15 (1988)

229-236 2. E. Reissner, Applied Mechanics Review, 38 (1985) 1453-1464 3. J.M. Whitney, Journal of Composite Materials, 6 (1972) 426-440 4. S.A. Ambartsumian, Izv., Otd. Tech. Nank. AN SSSR 5 (1958) 69-77 5. Z. Kaczkowski, Plates. Statical calculations Warsaw:Arkady(1968) 6. V. Panc, Theories of elactic plates, Prague: Academia (1975) 7. E. Reissner, International Journal of solids and Structures,

11 (1975) 569-573. 8. N.J~aganO,Journal of Composite Materials, 3 (1969) 398-411

9. S. Srinivas and A.K. Rao, International Journal of solids and structures, (1970) 1463-1481

10. P. Lory, Renault - Direction de la Recherche, Rapport interne, (1987).

11. Q. Li~ and M. Touratier, Rapport n0 1 RNUR - SERAM, 1987

FREE-EDGE STRESS SINGULARITY COMPUTATION

P. DESTUYNDER, Y. OUSSET'

ECAMP - Grande Vaie des Vignes - 92290 CHATENA Y-MALABRY - France 'CETfM - 52 Avenue Felix Lauat - 60300 SENLIS - France

ABSTRACT

For a multilayered plate, the stress field can become singular in the neighbourhood of the interlayer/free edge intersections. The calculation of such stress singularities involves the determination of both the order and the magnitude of the singularities. The aim of this work is to describe methods that allow to compute these two quantities. At first, the orders are computed using a semi- analyti­cal method and then, the intensity factors are expressed as the ratio of two energy terms and computed using finite element solu­tions of the free edge problem. Numerical results are presented.

INTRODUCTION

In the neighbourhood of a singular point the displacement field can be written as (Grisvard/l/)

where uS is the singular part of u of the form

(2)

and uR its regular part.

726

The orders a are computed using a semi-analytical method (Destuynder/2/) wRile the intensity factors Kn are computed using the dual singular functions concept (Destuynder & Ousset/3/).

I - SINGULARITY ORDERS

The asymptotic behaviour of the stress field near A (see Fig.l-(C)) is given by the non-zero solutions with finite strain energy of the following problem

~ 0i •• (S) = 0 in 0+ and 0 -

[Oi~ (S)] = [Si] 0 on E (3)

= 0 + and r1: o n on rL ia a

where i=1,2,3, a=1,2 and [ ] is the jump of any quantity through E. The problem (3) is not solved directly. It is transformed into

two equivalent sets of equations by using the complementary energy theorem and the Airy's stress functions (see /2/).

- Four equations of compatibility in 0+ and 0-. They are solved analytically. This permits to express the Airy's functions as an unknown linear combination of twelve known elementary functions.

- Six jump conditions through I. These conditions added to the six boundary conditions on rt and r1: permits to show that the unknown coefficients are the non-zero solutions of a complex valued homogeneous linear system. The seeked values of a n (2) are the roots of the determinant of this system. They are found numerically by combining a GAUSS' factorization with a 2D dichotomy.

II - INTENSITY FACTORS

2.1. The edge effect problem

Let oKL be the stress field solution of a Kirchhoff-Love plate problem. The stress vector t defined by

t· = - cJ<L na 1 la (4)

does not vanish locally on the lateral free edge rL' We have only

Ir ti dx 3 = 0 L

(5)

Accordingly (Destuynder/4/), near r , the KL stress field must be corrected by the solution of the so-ca!led edge effect problem.

727

°ia.a (uc ) = 0 in OI ;I=l •..• N

[u9) = [aD (uc ») 0 1

(6)

°ia (uc).na = ti

0i3 (uc) = 0

lim (uC • oc) 0 x 2 ~ + 00

equipped with the plane strain constitutive equations. We have seen that this problem admits singular solutions.

2.2. The dual singular functions

Let A be a singular point. The dual singular functions S* at A are the images. in the space of loadings. of the singular displace­ments Sn/3/. There are as much functions S* as functions Sn. say N • Any function S* satisfies the homogeneous problem associated to (g) (t = 0) and can be decomposed in the sum

S* = S*S + S*R (7)

wher;.e - S S is the singular part (square int*grable) computed in the same

way of the Sn. but with -1 < Re (an) < O. The NS singular func­tions shall be noted S n.

- S R i~ the regular part (finite strain energy). The S n functions do not satisfy the equilibrium equations (jump conditio~s) in oI(II~ e~cept for the ones including A. The as so­ciate4, S R. not*ed*S* (S n.). are then defined in order that the sums S • noted S (S n). satlsfy the homogeneous problem.

Practically. for numerical computations. the infinite*~trig is trun­cated by a boundary ac (see Fig.1-(b» so that the S (S ) are so­lutions of the following problem

*R -oia.a(S ) *n

°ia.a(S ) in CI

S*R = _S*n on ac *R = o. (S*n) rL °ia(S ) na na on la *R aD (S ) *n -aD (S ) on r±

*R [oi3(S ») _ *n - -[oi3(S ») on II

Where CI is part of 01 included between rL and ac.

2.3. Application to the computation of ~.

(8)

In 131. it is shown that the intensity factors ~ are obtained by solving the linear system

728

where ~ is given by

~ = J [n.o(S*m).Sn-n.o(Sn).S*m] d aco aco

and bm is given by either

b = J t S*m dx + J [n.o(S*m).uc- n.o(uc).S*m] d aco m rLnc o ' 3 aco

or

Moreover,

if a: is the singularity order of S*m, then a: - 0mn «n - if the an are single roots, then Amn = 0mn Ann - if S*m and Sn are not related to the s.ame singular point, then ~ O.

III - NUMERICAL EXEMPLES

3.1. Isotropic case

TIle present method was used to predict singular stresses in the bonded joints shown in Fig.2. The material characteristics were the following ~B = .35 ; ~M =.3 ; EB = 3400 MPa ; EM = 200 000 MPa and the magnitude of the applied load was 1 MPa. Fig.3 gives the pattern of the singular stresses near B for an overlap lenght of 50 mm and an adhesive tickness of .2mm.

3.2. Anisotropic case

Here, the method is applied to the determination of the singu­larity orders along a circular hole of a multilayered plate. The material characteristics were

Fig.4 gives the variation of a with respect to the polar angle for [9, 9 + ~] and [9, 9 + rr/4] layers.

CONCLUSION

We have presented a fast method in order to compute stress sin­gularities near the free edges of a multilayered plate (much less than 30 s on a VAX or a workstation). This method can be implemented in any finite element code as a post-processor.

(a)

X3

A

A

n

~XO

~

r L

W

A

( /

j !

,II

I

x, X

3

r~j

Q+

I

.. X

2

Q-

(c)

rC

Fig

.l

-G

eom

etri

cal

no

tati

on

s fo

r th

e

edg

e p

rob

lem

1 Q

0+

1

X

N

(b)

r+

I'

I8

IN-1

r-

-..J

N

to

730

7.5

5.

2.5

1.

O. -1.

2 -SUPPORT

E - E Xl It) - A '1..45 0 • " <'! B

~ORT -ADHESIVE - E

~ E It)

I--

45 mm 50 mm 45 mm

Fig . 2 - Geometrical characteristics of the single lap bonded joint.

.1

a22

all mm)( 10-3

. 25 .5 .75 1 . X, 0'2

Stresses in the adhesive along the interface versus the distance from B.

a 1 .

. 99

.98

731

---[O,O+71/4] --[0.O+71/2] Fig.4 - Singularity order along the hole versus the polar ang­

le.

REFERENCES

1. Grisvard P., Elliptic problems is nonsmooth domains, Pitmann (1985).

2. Destuynder P., Comptes-rendus a l'Academie des Sciences se­rie II, 302/6 (1986) 257-262, Paris (France).

3. Destuynder P. & Ousset Y., Report 107060/2, CETIM (1987), SenEs (France).

4. Destuynder P., Une theorie asymptotique des plaques minces en elasticite lineaire, Editions Masson (1986), Paris.

EVALUATION D'UN NOUVEL ELEMENT FINI POUR L'ANALYSE STATIQUE OU DYNAMIQUE

DES PLAQUES COMPOSITES

P. LARDEUR, J.L BATOZ

Universite de Technologie de Compiegne - Departement de Genie Mecanique Division MNM - BP 649 - 60206 COMPIEGNE - France

RESlNE

A finite element formulation is presented for static and vibration analysis of multilayer composite plates. We consider a mixed plate theory including transverse shear deformation and stresses. The proposed element is based on the introduction of transverse shear equations in a discrete way. The element has three nodes and nine degrees of freedom. Numerical applications dealing with three and nine layers and sandwich plates show that our model is eftective for a wide range of structures.

I - INTRODUCTION

Contrairement aux materiaux isotropes, les modules de rigidite au cisaillement des materiaux composites sont faibles par rapport aux modules de rigidite en traction-compression. La prise en compte systematique du cisaillement, et particuli~rement du cisaillement transversal (C.T.), est necessaire dans les structures de type poutre, plaque et coque. Les theories basees sur des hypotheses cinematiques sont generalement correctes pour I'eftet de flexion, mais pas pour Ie C.T .. Les conditions de continuite des contraintes de C.T. ne sont pas satisfaites aux interfaces. Nous avons retenu et adapte un mod~le mixte 11/, qui respecte toutes les conditions necessaires. En ce qui conceme Ie mod~le numerique, de nombreux elements finis de plaque avec eftet de C.T. ont deja ete developpes. En general, leur comportement est mauvais quand la plaque devient mince (blocage en cisaillement). Pour resoudre ce probl~me, de nombreuses methodes ont ete proposees : I'integration reduite ou selective, les formulations mixtes, I'utilisation

734

de variables supplementaires. Nous proposons ici un nouvel element fini de plaque a 3 noeuds et 3 ddl par noeud, base sur I'introduction sous forme discrete des hypotheses de plaque avec C.T.

II - MODELE THEORIQUE

2.1. Hypotheses cinematiques et statiques

Nous choisissons Ie champ de deplacements de Mindlin-Hencky-Uflyand, soit :

f U(x,y,z) = zPx (x,y)

\ V(x,y,z)=zPy(x,y)

~ W(x,y,z) = w(x,y)

( 1 )

ou Px et Py representent les rotations de la normale et w Ie deplacement transversal. Les contraintes planes (fxx, (fyy, (fxy deduites de ce champ, sont lineaires par couches et discontinues aux interfaces.

Les contraintes de C.T., (fxz et (fyz, sont calculees a partir des contraintes planes par integration des equations d'equilibre tridimensionnelles. Ces contraintes sont quadratiques par couche, et leur continuite est assuree en tout point de I'epaisseur.

2.2. Formulation variationnelle

Nous utilisons Ie principe mixte d'Hellinger-Reissner modifie tridimensionnel, soit :

avec <X> = <Px,x Py,y Px,y + Py,x> ; < 'Y> = <w,x + Px W,y + Py>

<t> = < O'xz (fyz>

[OJ (3 x 3) et [DcI (2 x 2) contiennent les coefficients elastiques en un

point dans Ie repere x, y, z.

Apres integration suivant z et minimisation de la fonctionnelle par rapport a <t>, on obUent la loi de comportement en cisaillement et on retrouve une fonctionnelle de type {mergie potentielle totale.

735

III - MODELE ELEMENTS FINIS

L'element appele DST est deerit de fa~n plus detaillee en [2] et [3]. Nous rappelons ici les principales etapes de sa formulation. En tenant compte des equations d'equilibre, et de la loi de comportement nous obtenons :

(1) = (F C) =(FX (Di ) ~x' ~y) ; FY(Di } ~x' ~y) ( 3)

ou FX et FY s'expriment en fonction des caracteristiques elastiques et des derivees secondes de ~x et ~y. Si la structure n'est pas sensible au C.T., Ie resultat du systeme (3) est <y> = <0> (modele de Kirchhoff). Dans I'element fini DST (Discrete Shear Triangle), les rotations ~x et ~y sont definies par des polynOmes

quadratiques incomplets a 9 termes (Le. ~s est quadratique et ~n est lineaire Ie long des cOtes des elements) et west cubique Ie long des cOtes ; ainsi FX et FY sont constants sur I'element. Les relations (3) sont introduites sous forme discrete sur Ie contour de I'element, sous une forme inspiree de celie utilisee dans I'element de plaque mince DKT 171, soit :

• <y> = <FC> aux noeuds sommets 1, 2, 3

• w's + ~s = -S FX + C FY aux points milieux 4, 5, 6

ou C et S sont les cosinus directeurs du cote (voir Fig. 1).

(4 )

( 5 )

Ces relations permettent d'eliminer 9 variables soit w,x ,W'y aux noeuds

sommets et ~s aux milieux des cotes. Les variables nodales de DST sont wi, ~xi, ~yi

aux troix noeuds sommets. L'energie interne de deformation d'un element est definie par:

(6 )

(7)

avec (un>= (Wi ~xi ~yi i=1,2,3)

L'element DST presente les avantages suivants :

nombre de ddl minimum, - pas de blocage (shear locking) dO au C.T., puisque sa matrice de rigidite

degenere naturellement vers celie de I'element DKT 171, les matrices [KF] et [KC] etant integrees exactement,le rang de [K] est systematiquement correct, les qualites de convergence vers les solutions analytiques sont toujours satisfaisantes,

736

- la procedure, appliquee a la poutre, donne la matrice de rigidite exacte des poutres avec C.T.,

Pour I'analyse dynamique, nous formulons la matrice masse, en utilisant les hypotheses (1), avec w, Px, Py lineaires sur I'element.

IV - APPUCATIONS NUMERIQUES

- Plaque carree a 3 couches 0/90/0 ou 9 couches 0/90/0/90/0/90/0/90/0 simplement supportee soumise a un chargement doublement sinusoIdal : la structure est constituee d'un materiau orthotrope (graphite-epoxy), dont les caracteristiques mecaniques sont donnees en /4/. La figure 2 montre I'influence du cisaillement transversal sur la fleche au milieu de la plaque. La figure 3 montre la distribution de crxz sur Ie bord de la structure. La figure 4 montre

I'influence du cisaillement transversal sur la premiere frequence propre. Nos resultats sont en bon accord avec les solutions de reference /4/, /5/.

- Plaque carree sandwich simplement supportee sou mise a un chargement uniforme p : la structure est constituee d'un materiau faiblement orthotrope (cristaux d'Aragonite), dont les proprietes mecaniques sont donnees en /6/. Le coefficient c determine Ie rapport de proportionnalite entre les proprietes mecaniques des peaux et du coeur : iI prend successivement les valeurs 1, 10, 50. L'analyse est faite pour un seul rapport Uh = 10. Dans Ie tableau 1, nous donnons les resultats concernant la fleche au centre de la plaque. Dans Ie tableau 2, nous donnons les frequences propres fondamentales. Plus c est grand, plus les resultats s'eloignent de la solution classique de Kirchhoff. Dans tous les cas, nos resultats sont en bon accord avec la solution de reference /6/.

V - CONCLUSION

Le nouvel element DST, base sur une tMorie de plaque mixte avec C.T., permet d'obtenir de bons resultats en statique et en dynamique pour les structures composites. Toutes les grandeurs importantes (deplacements, contraintes, frequences propres) sont obtenues avec une bonne preCision. Le modele est particulierement performant pour les structures a grand nombre de couches, et aussi pour les structures sandwich.

REFER8\CES

1. E. Reissner, AIAA Journal, vol. 10, n° 5 (1972) p. 716. 2. P. Lardeur et J.L. Batoz, AUM, Actes du 8eme Congres Francais de Mecanique,

Nantes, tome 2 (1987) p. 288. 3. J.L. Batoz et P. Lardeur, IJNME, a paraitre. 4. N.J. Pagano et S.J. Hatfield, AIAA Journal, vol. 10, n° 7 (1972) p. 931. 5. J.N. Reddy et T. Kuppusamy, J. of Sound and Vibration, vol. 94 (1984) p. 63. 6. S. Srinivas, Journal of Sound and Vibration, vol. 30, (1973) p. 495. 7. J.L. Batoz et K.J. Bathe et L.W Ho, IJNME, vol. 15 (1980) p. 1771.

c = 1 c = 10 c = 50

DST 6 x 6 180.88 41.92 16.65

Elasticite 141 181.05 41.91 16.75

DST 6 x 6 Uh = 10000 166.94 30.96 6.77 (= DKT)

TMorie plaque mince 141 168.38 31.24 6.76

Influence du C.T. + 7 % + 25.5 % + 59.6 %

Table.!u 1 - Fleche au centre d'une plaque sandwich­w = w max C3J 2<.2) I h P - Influence du C.T.

c = 1 c = 10 c = 50

DST6x6 0.0929 0.1920 0.3029

Elasticite 161 0.0925 0.1913 0.2995

DST 6 x 6 Uh = 10000 0.0978 0.2272 0.4864 (= DKT)

Theorie plaque mince 161 0.0969 0.2249 0.4795

Influence du C.T. - 4.8 % - 17.6 % - 60.1 %

Tableau 2 - Frequence ro re fondamentale d'une plaque sandwich -2 2

A. = P (2) h ro I D3S< 2) - Influence du C.T.

2

w

cote i-j

Fig. 1 - Presentation de I'element DST a 9 ddl

737

738

5

0 ... DST 6·6 9 couches

4

\~--<>- ELASTICITE 9 couches .- DST 6·6 3 couches

3 '0- ELASTlCITE 3 couches WIWK

.,6- KFlCHHOFF 2

A--A 2

O+--------+--------~-------+--------~------~ o 10 20 30 40 50

Llh

Fig, 2 - Plaque composite 3 ou 9 couches: influence du cisaillement transversal sur la fleche maximum

~:: 0

1_ -._--\.-- <\". '.- DST 6·6

0, 3 ~ ~ "'0 '0- ELASTICITE

0 , 2 ~o

~:~+------+------+------+I------~I~----~ 'O'O'tOO 0,05 0 ,1 0 0 ,1 5 50 ~O'2~1GXZ · 0 , 2 ~~

:~:: t-A/ •

· 0 ,50-

Fig, 3 - Plaque composite 9 couches: distribution de sigmaxz en (0, U2) pour Uh = 10

16~~~~~~~~==~~~ 14 ~

12

10

Cll 8

6

4

2

• DST 6·6 3 couches

o E. F. 3D 3 couches

• DST 6·6 9 couches

a~

O+---r---~~~~---+---+--~--~--~~

o 10 20 30 40 50 60 70 80 90 100 Llh

Fig. 4 - Plaque composite 3 ou 9 couches: influence du C.T. sur la frequence propre fondamentale

CISAILLEMENT SHEAR

Chairman: Dr I. VERPOEST Katholieke Universiteit Leuven

AN EXPERIMENT AL·ANAL YTICAL INVESTIGATION OF INTRALAMINAR

SHEAR PROPERTIES OF UNIDIRECTIONAL CFRP

w. BROUGHTON, M. KUMOSA, D. HULL

University of Cambridge - Department of Materials Science & Metallurgy Pembroke Street, CB2 3QZ CAMBRIDGE - England

ABSTRACT

Intralaminar shear properties of four 40-ply unidirectional carbon-fibre reinforced polymer matrix composites were investigated using the Iosipescu shear test. The apparent shear strength and shear moduli were measured using specimens with two different fibre orientations. Finite element analysis was applied to determine the stress distribution within the Iosipescu specimen. The experimental shear moduli when corrected to account for the non-uniform shear stress distribution in the gauge section, were within 8% of micro-mechanics predictions. The failure modes in the two fibre orientations were different, resulting in a change in apparent shear strength.

INTRODUCTION

Ideally, for quantitative intralaminar shear measurements, the shear test method should provide a region of pure, uniform shear stress. The test method should also be relatively simple to conduct, employ small, easily fabricated specimens, and be capable of measuring both shear strength and shear modulus. Adams and Walrath/1-2/ proposed that the Iosipescu shear test technique employing a double edge-notched flat laminate specimen would satisfactorily meet the above conditions.

Numerous numerical and experimental studies on the distribution of normal stresses a and a , and shear stress c within the Iosipescu specimen hav~ shownYthat the shear stres~Y distribution in the test section of both isotropic and orthotropic specimens is not uniform/2-S/. These stresses have been shown to

742

be a function of loading conditions, orthotropy ratio E11/E22 and notch geometry. Therefore, the average shear stress:

• = PiA (1)

where P is the applied force and A is the net cross-section area between the notch roots, is generally not representative of the shear stress at the specimen centre. The numerical studies indicate the need to account for the difference between the applied shear stress and the actual stress in the gauge section.

This paper presents a combined experimental and analytical investigation of the intra laminar shear properties of unidirectional carbon/epoxy and carbon/PEEK composites. The apparent shear strength and shear moduli were measured for specimens with two different fibre orientations. Finite element analysis was used to interpret the experimental results.

EXPERIMENTAL PROCEDURE

The intra laminar shear strength and shear modulus was measured for 40-ply unidirectional AS4/3501-6, XAS 914C and T300B R23 carbon/epoxy, and APC-2 carbon/PEEK using the Iosipescu shear test technique. Iosipescu specimens were cut with fibres parallel to the sgecimen longitudinal axis [00 ] and perpendicular to this axis [90]. The dimensions and fibre orientations of the specimens are shown in Fig.l. A 900 angle notch was cut with faces oriented at ± 450 to the fibre direction, to a depth of 20% of the specimen width. The notch root radius was - 40~m.

The specimens were tested at a cross-head speed of 1mm/min. A Kyowa (KFC-1-DI6-11) 900 biaxial rosette was bonded to the centre of the specimen in the area between the notches. The gauges were aligned at ± 450 to the specimen axis and were 1mm long. Tests were performed in a universally adaptable In-plane Biaxial Stress Fixture (Fig.2)/6/. This test fixture is capable of testing Iosipescu specimens in either shear or a combination of shear and transverse tension or compressive loadings.

Stress distribut~ons were determined by finite element methods using the PAFEC system. The loading boundary conditions applied to the finite element model were chosen to simulate the loading originally proposed by Iosipescu with two force couples appli~d as shown in Fig. 1 (force couple condition)/4/. A comparative photoelastic study on isotropic and orthotropic Iosipescu specimens was used to show that the force couple conditions were representative of the actual experimental configuration.

RESULTS AND DISCUSSION

In general, stress-strain linearity extended to stress levels between 40 and 50% of the apparent shear strength. Axial splitting was observed for [00 ] specimen/8/; manifested by two

743

successive load drops of the order of lOON in magnitude. These load drops corresponded to the initiation of two cracks immediately adjacent to the notch roots, on the opposite sides to the inner loading points. After crack formations, two damage zones at approximately 800 to the crack planes appear with barely detectable influence on the shear stress-strain response. This damage zone consists of short interfacial shear cracks, fibre bending and out-of-plane plastic deformation for both the carbon/epoxy and carbon/PEEK specimens. After the appearance of these failure modes, the material continued to bear load with an increased non-linearity. Failure in [900 ] specimens/8/ was catastropic and resulted in the complete separation of the specimen along the notch root axis.

The measured shear moduli and stresses at failure (L1~) are given in Table 1. Failure stresses correspond to the average shear stresses (equation 1) recorded at the onset of initial splitting at the notch roots for both types of specimens. Each data point is the average of 5 test results. The apparent shear moduli for [900 ] specimens (G12T ) are significantly smaller than the apparent shear moduli (G12L ) for [00 ] fibre orientation. From the finite element analysis/41 the correction factors were determined to account for the non-uniform shear stress distribution along the notch axis. The estimated factors were 0.84 and 1.13 for G12L and G12T respectively for E11/E22 = 14.2.

Predicted shear moduli were obtained according to the Halpin-Tsai equations/?/:

and G12 /Gm = (1 + 411Vf )/(l - ilVf )

II = (Gf/Gm-l)/(Gf/Gm + 1)

(2)

(3)

where Vf = volume fraction, Gf = fibre shear modulus, Gm = matrix shear modulus and the reinforcement factor E; 1. The corrected experimental results agree to within 8% of the predicted data (Table 1).

Kumosa and Hull/4/ showed that the shear stress concentration factor Kt at the notch roots, defined as the ratio of the shear stress at the notch root to the stress in the specimen centre, can be expressed in terms of the orthotropy ratio and the fibre orientation as:

(4)

where A = numerical isotropic factor dependent on the loading boundary conditions. Similarly, a generalised numerical relationship based on the finite element data/4/ can be used to correct G17 for the error introduced by the nonuniformity of the shear stress distribution along the notch root axis:

G1Z /G 1; = -B (E 11 /E 2Z )li4 + C (5)

744

* where G ~ is the apparent intralaminar shear modulus, Band Care constan!~ dependent on loading boundary conditions.

The comparative photoelastic and numerical studies revealed that the general features of the normalised von Mises stress contours for both [00 ] and [900 ] specimens showed good qualitative agreement with the corresponding isochromatic fringe patterns (Figs. 3 and 4)/8/.

The finite element analysis assumes a linear-elastic response of the Iosipescu specimen to failure, with predicted Kt values of 2.08 and 0.55 for [00 ] and [900 ] specimens respectively. In actual fact, the shear stress-strain response was non-linear and Kt approached unity at failure. The non-linear response is caused by creep, plasticity and accumulation of micro-damage at the notch roots/8/.

CONCLUSION

1. Using the correction factors from finite element analysis the actual shear moduli can be established from the apparent data for [0°] and [900 ] specimens.

2. A simple numerical relationship has been determined to establish the actual shear modulus.

3. Comparative photoelastic and numerical studies have confirmed that the force couple condition is very close to the experimental loading configuration.

ACKNOWLEDGEMENTS

The authors would like to thank the Aeronautical Research Laboratories, Victoria, Australia and the Science and Engineering Research Council, U.K. for supporting this study.

REFERENCES

1. Adams D.F. and Walrath D.E., Experimental Mechanics, 23 (1983) 105-110.

? Walrath D.E. and Adams D.F., Analysis of the Stress State in an Iosipescu Shear Test Specimen, University of Wyoming Report No. UWME-DR-301-102-1, June 1983.

3. Barnes J.A., Kumosa M. and Hull D., Composites Science and Technology, 28 (1987) 251-268.

4. Kumosa M. and Hull D., International Journal of Fracture, 35 (1987) 83-102.

5. Pindera M.J., Choksi G .• Hidde J.S. and Herakovich C.T .• Journal of Composite Materials, 21 (1987) 1164-1184.

6. Broughton W.R., Mixed-Mode Failure of Carbon-Fibre Reinforced Composite Materials. University of Cambridge, UK. Technical Report, June 1987.

7. Halpin J.C., Revised Primer on Composite Materials Analysis. (1980) Technomic Publishing Co.

8. Broughton W.R •• Kumosa M. and Hull D., to be published.

745

TABLE 1

COMPARISON OF APPARENT, CORRECTED AND PREDICTED SHEAR MODULI AND SHEAR STRENGTH

APF.ARENT CORRECTED PREDICTED APF.ARENT SPECIMEN G1;(GPa) G1/GPa) G12 (GPa) 't 1;(MPa)

AS4/3501-6[OoJ 7.48±0.05 6.28 6.15 69.12:t0.55 AS4/3501-6[90 1 5.87±0.20 6.63 58.69±1.46 XAS 914C [OoJ 6.3l-tO.25 5.30 5.70 75.11 t 1. 84 XAS 914C [90 ) 5.01±0.41 5.66 62.02±1.86 T300B R23 rooJ 5.38±0.23 4.52 4.66 69.11±2.11 T300B R23 [90 ) 4.44±0.27 5.02 58.51±2.53 APC-2 roQJ 6.48±0.55 5.75 5.46 71.52±0.41 APC-2 [90 1 5.05.1:0.19 5.71 71. OUL 29

a = 70nun b = 18nun w = 12nun

l!! y l Pa L = 80mm

a-b

x

L _ a

Fig. 1 - Schematic of Iosipescu Shear Test Specimen

Fig. 2 - In-Plane Biaxial Stress Fixture

746

Fig. 3

Fig. 4

~· ; ~ " fil'-" '~. .. ~ .. . -' . . . . ..'.. ")' ., ' . .

, " :. \1'.' " - . . , - '

~,.";, ,; / ( .. ~: . ' .~ ~ .,.

:' .• : . • r.~ I '~.> ... '.: .. , _ _ -,--, ' . ' ........... jJ---.-"-- _

[0°] Iosipescu Specimen Left - Isochromatic Fringes Right - Normalised von Mises Contours

- [90°] Left Right

Iosipescu Specimen - Isochromatic Fringes - Normalised von Mises Contours

INTERLAMINAR FRACTURE TESTING OF CARBON FIBRE/PEEK COMPOSITES

VALIDITY AND APPLICATIONS

P. DAVIES, W. CANTWELL, H. RICHARD, C. MOULIN, H.H. KAUSCH

Ecole Poly technique Federale de Lausanne, Laboratoire de Polymeres 32 Chemin de Bellerive- 1007 LAUSANNE - Suisse

ABSTRACT This paper examines the application of interlaminar fracture testing to carbon fibre reinforced PEEK (poly-ether-ether-ketone) composites An experimental study of over 100 specimens is reported, in which the influence of specimen thickness, stiffness, width and defect type have been systematically varied. This has revealed a significant effect of specimen stiffness in certain cases, which must be taken into account if useful results are to be obtained. Following the optimisation of the specimen the influences of fibre type (AS4 or IM6) and rate of cooling after moulding (50°C/min. down to 0.3 °C/min.) were studied.

INTRODUCTION The use of interlaminar fracture tests to characterize the delamination

resistance of carbon/epoxy composites has become widely accepted in recent years. Although no standard test methods exist yet, the values of GIc and GUc obtained from tests on unidirectional specimens of these relatively brittle materials are normally below l000J/m2 and appear reasonably reproducible. The relevance of these values to mixed mode loading of multidirectional laminates is less clear, but at present their principal application is in the assessment of new materials.

With the development of composites based on ductile thermoplastic matrices values of GIc of over 2000 J/m2 are now commonly quoted. However, due to the recent introduction of these materials little work has been carried out so far to investigate the specimen independence and validity of such tests.

The present paper describes an extensive programme of tests designed to optimise the DCB and ENF specimen geometries for testing of a carbonlPEEK

748

composite, and then two examples of their application are presented. The following test and material parameters were varied for both mode I and mode II ;

- specimen thickness (2, 3, and 5 mm) - specimen width (10, 15, and 20 mm) - starter defect thickness (12.5, 25, and 50~) - distance between supports, mode II (75,80 and lOOmm) - fibre type (lM6 and AS4)

- cooling rate after moulding (OJ, 1,3, and 50 °C/minute).

All results are presented in tabular form, but for clarity Figures 1 and 2 also highlight the influence of specimen thickness and cooling rate.

lEST METIIODS The mode I test specimen was the double cantilever beam (DCB), the use

of which has recently been reviewed in detail elsewhere /1/. Gle values were determined via an experimental compliance calibration as described by de Charentenay et al /2/. Specimens were loaded at 2mm/min.

Mode II tests were performed on edge notch flexure (ENF) specimens. This specimen has also been extensively studied, /e.g. 3-5/. The length between supports was 80mm unless specified, loading rate was 1 mm/min., and a folded PTFE film was inserted to reduce friction between the crack faces. The analysis of data from these tests was complicated by non-linear load-displacement plots. The approach used was therefore to calculate a minimum value of GIIc at the point of deviation from linearity using a beam theory expression /3/. A second value, corresponding to the maximum load, (an easier point to define accurately), was also determined as follows; an effective modulus was calculated from the initial slope and initial crack length. This modulus enabled an effective crack length to be determined corresponding to the compliance at maximum load, as if all the non-linearity was due to crack propagation, and then the value of GIIc at that point was calculated using this crack length.

All tests were performed on unidirectional material, and in each case at least 4 specimens were tested.

I INH.UENCE OF SPECIMEN VARIABLES IN MODE I.

1.1 THICKNESS Published results for propagation in AS4/PEEK show considerable

scatter, and a plot of these results as a function of thickness suggests that a dependence of Gle on thickness may be responsible for part of this scatter, Figure 1 /6/. Tests were therefore performed on specimens of 3 thicknesses, (2, 3 and 5mm) of IM6/pEEK and the results are also shown in Table 1 and Figure 1. The 5mm thick specimens give a mean propagation value of 3320 1/m2, a value identical to that quoted in the manufacturers data sheet, (which was also obtained using 5mm thick specimens.) The thinner specimens give much lower values,

749

around 2100 11m2, and this difference is attributed to multiple cracking occurring in the 5mm specimens 16/. Sections through the 2 and 3mm DCB's showed no multiple cracking, and little fibre bridging was observed during tests on these specimens.Values for the 2mm thick specimens have been corrected for large displacements n I.

A supplementary set of tests was also performed on 3mm specimens milled down from a 5mm thick panel and these results (Table 1) show that the thickness dependence of GIc values is not due to differences in the moulding of panels of different thicknesses.

1.2 WIDTH Specimens of three widths were tested, 10, 15 and the standard 20mm. A

small increase in GIc values was noted for thinner specimens, Table 2, which may be due to plane stress or edge effects. Indeed, sections through the thinner specimens revealed cracks on more than one plane near the specimen edges. Specimens of 20mm width appear to be the minimum necessary to avoid multiple cracking for 3mm thick specimens of this material.

1.3 TYPE OF DEFECT As the propagation mechanisms such as fibre bridging and multiple

cracking depend on specimen thickness it may be more desirable to quote initiation values of GIc, if these can be reliably measured. A study of the influence of the type of defect on initiation values has been described in more detail elsewhere, 18/, and Table 3 summarises the results obtained. A distinct rise in GIc values due to crack tip blunting is noted for 50J.1.m thick defects. In the tests described hereafter 25J.1.m folded aluminium foil defects were used.

1.4 A MODE I TEST METHODOLOGY From the results described above it is clear that great care is needed in

measuring GIc values in IM6/PEEK composites. Values from 1000 to 3500 11m2 have been obtained, in a single laboratory where all material has been made in­house from a single batch of prepreg, and a single method of data analysis has been followed. No discussion of data analysis methods is included here as such considerations have been treated elsewhere 11,9,101. The most pragmatic approach to characterizing the mode I delamination resistance of this material is simply to quote the lowest value obtained, but this is extremely conservative and seriously underestimates the toughness of this type of composite. Alternative approaches are being addressed in a number of task groups currently attempting to standardise the mode I test.

II INFLUENCE OF SPECIMEN VARIABLES IN MODE II

2.1 THICKNESS Values of GUc were obtained for the same range of thicknesses as in mode

I and the results are shown in Table 1. Once again larger values are observed for

750

5mm thick specimens, but it is believed that this is simply due to the quality of the mode I precrack. The difficulty in obtaining single-plane mode I propagation in these thick specimens is reflected in the subsequent mode II behaviour by a bunting effect at the crack tip, and this emphasizes the unsatisfactory nature of this approach.

2.2 WIDTH The same range of widths was used in mode II as in mode I, i.e. 10, 15,

and 2Omm. Table 2 shows the values of GIIc obtained, and a small effect of width is noted. Slightly higher values are noted for narrower specimens.

2.4 DISTANCE BETWEEN SUPPORTS Within experimental scatter no influence of distance between supports

was noted for the three distances tested, Table 4.

2.3 TYPE OF DEFECT The measurement of GIIc from a starter film gives artificially high values

/11/. The standard approach adopted for all the ENF tests presented here was to propagate a short mode I precrack, as this provides an easily-measurable initial crack. This technique is questionable, as has been shown in 2.1, as the development of multiple cracking and a bridged zone in 5mm thick mode I specimens will result in an unsatisfactory precrack. It is encouraging to note however, that for 3mm thick specimens, tests using a range of mode I precracks from 1 to 20mm did not influence subsequent GIIc values. In addition, a few tests in which mode II precracks were used gave consistently higher GIIc values for 3mm thick specimens, so there is some justification in taking the lower values measured from mode I precracks.

2.4 MODE II TESTING OF IM6/pEEK COMPOSITES. The determination of a value of GIIc for this material is complicated by

similar considerations to those met in mode I and once again values in the range from 1000 to 3500 J/m2 have been obtained. The influence of the type of defect or precrack, the detection of initiation, the treatment of non-linear behaviour and the data analysis methods all require further study before values can be used with any confidence.

III INFLUENCE OF FIBRE TYPE: AS4 v IM6 In order to compare the results from IM6/pEEK composites with those

for AS4 based materials, specimens of the latter of 3mm thickness were moulded. Comparative results are shown in Figure 1 and Table 1, and these indicate that at equivalent 3mm thicknesses the IM6 composites give slightly higher values. It should be noted that the values for 2mm thick IM6/PEEK specimens are also higher than those for the 3mm AS4/PEEK, as it might be suggested that equivalent specimen stiffness rather than thickness is required to compare materials whose

751

fibre modulus varies as much as those of IM6 and AS4, (278 and 235 GPa respectively).

The influence of fibre type cannot be dissociated from the changes in matrix behaviour due to the different fibre, and these are considered below.

IV INFLUENCE OF COOLING RATE. The cooling rate after moulding determines the degree of crystallinity

and the morphology of the PEEK matrix. Previous work has indicated a drop in mode I and mode II toughness in passing from material cooled at the

manufacturers' recommended rate (>10 °C/min.) to rates slower than 3 °C/min. /11/. In neither of the materials studied here, laid up from commercial production prepreg, was any significant evidence of sensitivity of Gc to cooling rate observed

at rates down to 1 DC/min., Figure 2. However, the AS4/PEEK specimens which were cooled very slowly, at 0.3 DC/min., showed very unstable propagation in mode I. Values are presented for stable propagation (a mean of 27 values from 3 of the 5 specimens tested), and also for the unstable propagation as characterized by the onset of instability and the arrest. There is a small drop in the value of stable GIc but the value of GIIc is not affected. The slowest- cooled IM6/PEEK showed no loss of toughness in mode I or in mode II.

DSC measurements in Table 5 show that the overall degree of crystallinity increases at slower cooling rates and that the fibre type does not affect this parameter. However, different structures may have the same overall degree of crystallinity, and the shapes and positions of the melting endotherms do differ between the two materials. These suggest more perfect, higher melting point crystal structure in the slowest-cooled (0.3 DC/min.) AS4/PEEK than in the equivalent IM6/pEEK composites. In further work morphological differences between these materials and those tested in the previous study (11) are being investigated, but the initial conclusion is that the mode I and mode II toughness of current production material are even less sensitive to cooling rate than previous tests had suggested . Cooling rate effects due to internal stresses may affect the fracture properties of multidirectional laminates however, and these are the subject of a separate study reported elsewhere /12/.

V CONCLUSIONS In this paper the influence of a number of parameters on the values

determined in mode I and mode II tests has been established. This has enabled appropriate specimen geometries to be selected (3mm thick and 20mm wide), to test two carbon/PEEK composites.

The influence of cooling rate after moulding has then been examined. The delamination resistance of IM6/PEEK was found to be insensitive to cooling rate while for AS4/PEEK extensive unstable mode I crack propagation occurred only in material cooled very slowly, at 0.3 DC/minute.

752

ACKNOWLEDGEMENTS This work was performed as part of a project financed by the Swiss "Fonds National", under programme NFPI9. Mr. Richard was funded by Rhone Poulenc, France, and the gift of AS4/pEEK prepreg from Dr. Carlile of ICI, Wilton is gratefully acknowledged.

REFERENCES 1. Davies P & Benzeggagh ML, "Interlaminar mode I fracture testing", to

appear in "Application of Fracture Mechanics to Composite Materials", Ed. K. Friedrich, Elsevier, (1989).

2. de Charentenay FX, Harry J-M, Prel YJ, Benzeggagh ML, ASTM STP 836, (1984), p84.

3. Russell AJ, Street KN, Proc. ICCM4. Tokyo (1984), p279. 4. Murri GB, O'Brien TK, Proc. 26th AIAA/ASME/ASCE/AHS Conf.,April

(1985), Orlando, Florida. 5. Carlsson LA, Gillespie JW, Tretheway BR, J. Reinf. Plastics & Composites,

Vol. 5, July (1986), p170. 6. Davies P, Cantwell W, Moulin C, Kausch HH, "A study of the delamination

resistance of IM6/PEEK composites", submitted to Composites Science & Technology.

7. Williams JG, Proc. ICCM6/ECCM2, London (1987), p3.233. 8. Davies P, Cantwell W, Kausch HR, "Measurement of initiation values of GIc

in IM6/pEEK composites", accepted for publication in Composites Science & Technology.

9. Whitney JM, Browning CE, Hoogsteden WJ, J. Reinf. Plastics and Composites, 1, (i98) p297.

10. Gu6dra D, Lang D, Rouchon J, Marais C, Sigety P, Proc. ICCM6/ECCM2, London (1987), p3.346.

11. Curtis PT, Davies P, Partridge IK, Sainty J-P, Proc. ICCM6/ECCM2, London, (1987) p4.401

12. Cantwell W, Davies P, Kausch HH, "The influence of cooling rate on the deformation and fracture of IM6/pEEK composites", submitted to J. Mat Science.

753

Table 1. SPECIMEN THICKNESS, 2h, (B = 20mm)

Material 2h (mm) GIc Prop. GUc (NL) GUc (Max)

IM6tpEEK 2 2140 (230) 1627 (192) 2506 (255) 3 2110 (320) 1390 (132) 2724 (264) 3* 2540 (230) ---- ----

5 3240 (370) 2344 (324) 3627 (251)

AS4tpEEK 3 1694 (131) 1095 (177) 2138 (318) ( * 3mm specimens milled down from 5mm panel)

Figure 1. Influence of specimen thickness on propagation values of GIc. Published values'for AS4tpEEK compared with values obtained in this study.

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754

Table 2. SPECIMEN WIDTH, B, (2h = 3mm)

Material B (nun) GIcProP· GUc (NL) GUc (Max)

IM6/PEEK 20 2110 (320) 1390 (132) 2724 (264) 15 2746 (120) 1367 (218) 2996 (255) 10 2516 (360) 1506 (252) 3100 (373)

Table 3. STARTER FILM THICKNESS, t, (2h = 3mm, B = 20mm )

Material t (J.Un) GIc Initiation

IM6/PEEK 12.5 1129 (63) 25 1270 (150) 50 1996 (496)

Table 4. DISTANCE BETWEEN SUPPORTS, MODE II ENF. (2h = 3mm, B = 20mm)

Material 2L(mm) GUc (NL) GUc(Max)

IM6/PEEK 75 1371 (168) 2684 (234) 80 1390 (132) 2724 (264) 100 1490 (125) 2657 (115)

Tables 1 -5 present mean values of Gc in 11m2, with standard deviations in brackets.

755

Table 5. COOLING RATE EFFECTS 9) = Degree of crystallinity, % by DSC, (2h = 3mm, B = 20mm, t = 251-1m)

Material °C/min 9) GIc Prop. GUc (NL) GndMax)

IM6/PEEK 50 25 (3) 2110 (320) 1390 (132) 2724 (264) 3 31 (2) 2363 (204) 1039 (234) 2523 (322) 1 36(1) 2480 (115) 1384 (137) 2870 (253) OJ 40 (4) 2212 (116) 1368 (73) 2398 (211)

AS4/PEEK 50 27 (2) 1694 (131) 1095 (177) 2183 (138) 3 33 (4) 1672 ( 64) 854 (161) 2387 (158) 1 36 (3) 1721 (149) 1147 (153) 2170 (233) OJ+ 39 (2) 1495 (142) 1091 (308) 2256 (140)

1583 (95)-1 725 (89)-A

+ All of these specimens showed very unstable "stick-slip" mode 1 propagation ..

Figure 2. Influence of cooling rate on values of Gc, for AS4/PEEK (.) and IM6/PEEK, (_)., 1 = Onset of Instability, A = Arrest

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METHODES NON DESTRUCTIVES NON DESTRUCTIVE TECHNIQUES

Chairman: Dr A. SAVADORI Enichem

MEASURING STRAIN IN CARBON FIBRE COMPOSITE LAMINATES USING THE RAMAN

OPTOMECHANICAL STRAIN GAUGE

F. UNDERWOOD, D. SHARPE', D. BATCHELDER"

British Aerospace Pic - Military Aircraft Division Richmond Road - KT2 5QS KINGSTON-ON-THAMES - England

"Aeronautical Engineering Department""Physics Department Queen Mary College - Mile End Road - E1 4NS LONDON - England

ABSTRACT

Polydiacetylene fibres were placed within a carbon fibre composite laminate in the ply direction during the lay-up procedure. Laser light was guided to the polydiacetylene (polyDCH) by means of a fibre optic. After curing the laminates were tested in four-point bending and the internal ply strains measured. The results showed that strains could be measured at any specific point within the laminate, they were consistent and the technique proved to be repeatable.

INTRODUCTION

As composites are being used more extensively, especially in the aerospace industry, the need for accurate strain measurement of laminated structures is obviously very important. The standard method of measuring strain is by attaching strain gauges to the surface of components, but the emphasis has now switched to determining internal ply strains.

A new type of strain gauge is required and research has centred on the use of fibre optics that utilise the elastic properties of the fibre optic to sense distributed strain [1-3].

This paper examines the development of a new technique that determines point strain within a

760

laminate. The Raman optomechanical strain gauge works on the principle that when laser light of a particular frequency is directed at a polydiacetylene fibre, the molecules within the fibre scatter some of the light inelastically. The scattered light has frequencies equal to the vibrational frequencies of the molecule and these frequency shifts are detected with a Raman spectrometer. On straining a polydiacetylene fibre the Raman band shifts linearly to a lower frequency and this shift has been calibrated for various polydiacetylene fibres. PolyDCH, a polydiacetylene stable at the high temperatures required for curing carbon _~ibre composites, has a frequency shift of -19.7cm /% strain [4-7] which is linear and elastic.

EXPERIMENTAL PROCEDURE

Fibre Preparation The sensor preparation began with the growth of needle like crystals of the diacetylene DCH monomer, 1,6-(N-carbazolyl)-2,4-Hexadiyne (DCH) [8]. These were then polymerised to polyDCH by exposure to 40 Mrads Co 'l-rays. The polyDCH fibres were checked under the microscope to ensure there were no visible defects such as cracking, twinning or' pi tting' • Fibres with a diameter of approximately 100/lm were selected, so as to match the core diameter of the optical fibre, and with lengths of between 8mm and 12mm.

Raman Spectroscopy

The resonance Raman spectra were taken with the output of a 10mW HeNe laser. The Raman microscope (shown in figure 1, the complete experimental set-up) was used to focus the laser beam on the end of the fibre optic and deliver the 180 oback-scattered light to the spectrometer. The spectra were dispersed across a Wright Instruments CCD (charge cooled device ) detector to permit -1 rapid and accurate data collection. The 2085cm Raman band associated with the vibration of the triple bond on the backbone of the polyDCH was used to measure the fibre strain.

Embedded Fibre Experiments

Laminates of dimensions 100mm x 25mm x 1. 016mm were made using Ciba-Geigy IM6-6376, XAS-914c and XAS-913c materials. The polyDCH fibre was placed between two 0° plies at O. 381mm from the neutral axis 100/140/lm graded indexed multimode fibre optics were chosen which optimised the core diameter to outer diameter ratio.

761

Figure 2 shows how the DCH and fibre optic were incorporated into the laminate. The polyDCR was placed along the 0° ply and the fibre optic butted up at right angles to it. A small amount of adhesive (Ciba-Geigy AYll1/RY111) was applied to 'spot bond' the joint. The cured laminates were then tested in four-point bending and Raman spectra were taken for each loading position. Laminates with the following lay-ups have been used :­a) (O,0,±45)s, the OCR placed between plies 1 and 2, the two 0° plies. b) (O,O,O,O)s, the DCH placed between plies 1 and 2.

A number of experiments were carried out to determine the effects of changing various parameters. i) Curing temperatures. Laminates of XAS-913c (90°curing temperature), IM6-6376 and XAS914c (175° curing temperature) were compared. ii) DCH fibre length. Fibres ranging from 8mm to 12mm were embedded in four different laminates. iii) Lay-ups. The strains in the laminates of the lay-ups above were compared. iv) The use of AY111/HY111 adhesive to spot bond the fibre optic/DCH bond. Laminates were laid up with the fibre optic tip released, the centre of the DCH fibre released and using only AY111 (no hardener). The resul ts from these three laminates were compared to those laminates cured using the 'spot bonding' technique.

RESULTS AND DISCUSSION

The theoretical strain shown in the following results has been calculated from the four-point bending geometry assuming simple beam theory. As can be seen in the resul ts the measured strain (which is calculated using !qe calibration for a free standing fibre of -19.7 cm /% strain) is considerably lower than the theoretically predicted strain. The reasons for this are being investigated and possible explanations are:-i) by using four-point bending the OCR is also put into bending and the effects of this are unknown; ii) the intrusion of the fibre optic within the laminate may be disturbing the strain distribution around the OCR fibre. To determine whether this is the cause of the low strain dependence micrographs have been taken and a 3-dimensional finite element analysis is being carried out.

The results presented in figure 3 show that the curing temperature does have an effect on the measured strain dependence. These results indicate that thermal

762

stresses are built up in the direction of the fibres during the curing process.

Figure 4 shows that all four laminates tested with differing embedded DCH fibre lengths exhibit the same strain dependence. For fibres of lOO~m diameter the stress transfer length is ~ lmm, which reduces as the volume fraction increases. The lengths of the DCH fibres ensure that the strains are measured at a point outside the region of stress transfer.

The two different lay-ups tested have no effect on the strain dependence as can be seen in figure 5. This is as predicted.

The effect of using a compatible adhesive to spot bond the fibre optic/DCH joint is shown in figure 6. Releasing the optic, the centre of the DCH fibre and using an adhesive that will not cure shows that the strain dependence is unaltered.

CONCLUSION

The results show that the Raman optomechanical strain gauge has been successfully adapted for use in composite laminates. Although lower than the theoretically predicted strains, the measured strains are linearly dependent on the applied load and are consistent for a given curing temperature.

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ACKNOWLEDGEMENTS

1.6 2 2.6 3 3.6 4 'rioo\ llIIIp!e_" <-J

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This research was supported by grants from the Science and Engineering Research Council and British Aerospace pIc.

REFERENCES

1. E.Snitzner, J.R.Dunphy and G.Meltz, "Fiber optic strain sensors", Proc. 1st International conf. on Fiber Optic Rotation and Related Technologies, (1982, Springer, New York). 2. C.D.Butter and G.B.Hocker, "Fiber Optic strain gauge", Applied optics, Vol 17, No.18, (Sept. 1978). 3. K.D.Bennett and R.O.Claus, "Microbending losses of optical fibers in composite materials", Proc. Optical Society of America Meeting (Washington DC), (Oct 1985) . 4. D.N.Batchelder,D.Bloor,"Strain dependence on the vibrational modes of a diacetylene crystal", J. of Polymer Sci.: Poly. Phys. Ed., Vol 17, (1976), 569. 5. C.Galiotis, R.J.Young, P.Yeung, D.N.Batchelder, "The study of model polydiacetylene/epoxy composites", J. of Mtls. Sci., Vo119, (1984), 3640. 6. C.Galiotis, R.J.Young, D.N.Batchelder, "A Resonance Raman Spectroscopic study of the strength of the bonding between an epoxy resin and a polydiacetylene fibre", J. of Mtls. Sci. Letters, Vol. 2, (1983),263. 7.I.M.Robinson,R.J.Young, C.Galiotis, D.N.Batchelder, "Study of model polydiacetylene/epoxy composites. Part 2 Effect of Resin shrinkage", J. of Mtls. Sci., Vol22, (1987), 3642. 8.C.Galiotis, "Polymer Single Crystals", PhD Thesis, University of London, (1982).

RAMAN OPTOMECHANICAL STUDIES ON FIBRES AND COMPOSITES

C. GALIOTIS, H. JAHANKHAN, I. MELANITIS, D. BATCHELDER

ABSTRACT

Queen Mary College - Materials Department Mile End Road - E1 4NS LONDON - England

The technique of Laser Raman Spectroscopy (LRS) applied in the study of fibres, composite interfaces is reviewed in this paper. An optical mechanical strain gauge has been developed for the determination of the internal strain in composite materials. The strain sensitive property of the high performance fibres is a vibrational frequency which may be determined by LRS. Examples are given for applications such as fibre characterisation, strain mapping in composites and examination of the strength of fibre/ matrix interfaces.

INTRODUCTION

The dependence of the molecular vibrational frequencies of crystalline fibres upon applied load can be very important in composite applications since the stress or strain that the reinforcing fibres experience can be monitored by measuring their Raman or IR frequencies /1/. In this context spectroscopic techniques which have been used in the past to characterize materials can also be used in a non-destructive way to measure directly stress or strain in the reinforcing fibres at the microscopic level.

In this paper the use of Laser Raman Spectroscopy applied in the characterization of carbon, as well as, aramid (Kevlar 49) fibres and on fibre/matrix interfaces will be reviewed.

1. CHARACTERIZATION OF FIBRES 1.1 Raman spectra

Carbon fibres are known to contain microcracks, voids, impurities

766

etc. on their surfaces all of which have an effect on the interfacial adhesion and in the mechanical properties of carbon fibre composites as a whole /2/. Raman spectroscopy has been demonstrated to be a useful technique for the characterization of carbon fibre surfaces due to the fact that the laser penetration depth is only approximately 500nm if the 514nm line of an argon ion laser is used /3/.

Figure I shows the Raman spectrum of a high modulus (HM) carbon fibre at the first order region. The band at approximately 1580cm-1 can be identified with the deformation mode of the hexagonal structure of the benzene ring and has been observed in graphite. The band at 136lcm-1 is not present in Raman spectra of single crystal of graphite but appears in the spectra of all disordered forms of carbon. The appearance of this forbidden band in the Raman spectra of carbon fibres has been attributed /3/ to the breakdown of translational syuunetry caused by the presence of microcrystalline structure on the surface of the fibres. Therefore this band can be considered as a "disordered-indicator" band and in fact the ratio of the l36lcm-1 to the l580cm- 1 gives a measure of the disorder present on the surface of the fibres regardless of the experimental conditions etc. Tuinstra and Koenig /3/ have related the ratio of these two bands to crystallite size by performing Raman as well as X-ray experiments on carbon fibres and polycrystalline carbon powders.

It has been found in this work that the relative intensities of the 1360cm-l and l580cm- 1 bands are particularly sensitive to surface treatment. As is shown in Table 1 for fibres of similar Raman spectra the normalised intensity of the IX (treated) band has twice the value of the normalised intensity of the OX (untreated) fibres. This indicates that the applied surface treatment on the fibres reduces dramatically the surface crystallinity and this possibly contributes to the two-fold increase in fibre/matrix adhesion for the IX fibres, as will be shown in a future publication.

1.2 Strain dependence of Raman frequencies Raman spectroscopy can be also used to measure the response of a

crystalline fibre to applied stress at the molecular level. As the frequencies of the Raman bands are directly related to the interatomic force cons~ants these shifts provide a clear measure of the way in which stress is distributed within the molecular structure of the fibre. This has already been demonstrated in the case of polydiacetylene /4/, poly( p-phenylene terepthalamide ) (Kevlar 49, Twaron) /5/ , carbon /6/, as well as, poly (p-phenylene benzobisthiazole) (PBT) fibres /7/.

Figure 2 gives the strain dependence of the two Raman bands for a HM carbon fibre. Numerical values for the rates of frequency shift are given in Table ,1. As can be seen from Figure 2 the relationship is linear up to fracture while the curves for loading and unloading are identical. Since Raman spectroscopy is only sensitive to bond extension two conclusions can be drawn : (a) the applied load contributes uniformly to bond extension, hence the linearity of the frequency vs strain curve and (b) no irreversible effects occur under cyclic deformation.

Kevlar 49 fibres, however, exhibit, a different behaviour as can be seen in Figure 3. Upon application of load the bonds/crystallites are extended uniformly and the relationship between frequency and strain

767

is linear but upon unloading a greater percentage of the load is now carried by the bonds I crystallites as indicated by the increase of the frequency vs strain slope (Figure 3). Indeed it has been reported elsewhere 181 that a strain hardening effect is observed upon deformation of Kevlar 49 leading to fibres of higher molecular alignment.

2. INTERFACIAL SHEAR STRENGTH (ISS) MEASUREMENTS

2.1 The Raman technique The relationship between Raman frequency and strain shown in

Figures 2 and 3 enables us to measure the applied load on a fibre in an area as small as 1 m by simply recording the Raman frequency at that point. In this way the fibres are converted to strain gauge sensors within a composite- provided that the matrix is reasonably transparent- with the spectrometer being the remote detector.

As far as ISS measurements are concerned it is clear that the values of t obtained from the existing methods depend greatly on the geometry as well as the fabrication conditions of the test. The advantages that the Raman technique can offer over existing techniques are the following: (a) the transfer lengths can be measured at any level of applied strain from the ends of a short fibre (b) in the case of a long fibre specimen the transfer lengths can be measured at a level of applied load sufficient enough to induce a single crack (c)fragmentation processes on single filaments can be monitored (d) the strength of the interface can be assessed under cyclic deformation (fatigue testing) (e) the test is valid not only for loading under tension but also under compression or other complex loading conditions and. finally. (g) it can applied not only to single fibre model composites but also to full composites.

The results obtained from a short Kevlar 491 Ciba Geigy 1927 system which has been recycled three times to 1.6% in tension are shown in Figure 4. The strain data have been normalised to the maximum strain reached at the middle of the fibre. It can be clearly seen that at 0.5% of applied strain (cycle 1) the load is built up from the ends of the fibre and becomes maximum at the middle. as predicted from the shear lag model 191 which assumes perfect adhesion between the two constituents. However at 1.6% of applied strain -end of the third cycle­the stress transfer profile degenerates to an approximately straight line and the transfer length increases dramatically. These observations indicate that the load is now transferred only by friction and therefore the chemical bonding between fibre and matrix has been severed. By employing an analytical model of an isolated fibre into an unbound matrix 1101 an interfacial shear strength t. of 55MPa has been measured The value obtained from the slope of the curve of Figure 4 at 1.6% applied strain is only llMPa • therefore. it is concluded that the interfacial adhesion is reduced by a factor of five by subjecting the model composite to cyclic deformation.

Fragmentation tests are also conducted in our laboratory and the transfer lengths. load transfer profiles. as well as. the overall fragmentation process at each level of applied strain are monitored using the Raman technique. Preliminary results from an HMS4 (Hercules) carbon fibrel Ciba Geigy 1927 epoxy system are shown in Figures Sa

768

and 5b for strain levels of 1.0\ and 1.7\ respectively. The overall load transfer profile at 1\ strain (Figure Sa) indicates that the load is built smoothly from the cracks on each side of the fragment and becomes maximum at the middle. It is interesting to note that considerable debonding is initiated at the vicinity of the crack tip even at a strain as low as 1\. The load transfer profile of a fragment at 1.7\ strain (Figure 5b) shows that (a) debonding near the crack tip can propagate dramatically along the fibre as the applied load increases and (b) the stress transfer profile is changing shape towards the linear built up observed in the weakly bonded situation of Figure 4. The present data show clearly that the "fragmentation test" as described above (section 3.Z) cannot provide an accurate method for the determination of the interfacial shear strength as it largely ignores the propagation of debonding from the crack tips, as well as, the physicochemical changes occurring at the fibre/matrix interface at the high strains required for full fragmentation. As it has already been demonstrated the Raman technique presents the only viable alternative for monitoring fragmentation processes and conducting interfacial shear strength measurements. In a future publication /11/ a complete strain mapping of the fragmentation process, as well as, the variation of the interfacial shear stress with applied strain will be presented.

CONCLUSIONS

It has .been demonstrated that Laser Raman spectroscopy can be used successfully to (a) characterize high performance fibres such as carbon or Kevlar and (b) determine the interfacial shear strength between a fibre and a polymer matrix. In fact, the Raman method appears to be far superior than any other existing non-destructive technique in the field of composites, since it can determine quantitatively the level of stress or strain in the reinforcing fibres.

ACKNOWLEDGEMENTS Courtaulds pIc. and Ciba Geigy are thanked for suppling materials

for this work.

REFERENCES 1- R. P. Wool and S. R. Bretzlaff (1986):J.of Pol.Sci.: PartB:

Pol.Phys.Physics edit., vol. Z4, pp.1039-1066 Z- I. L. Kalnin and A. Jager (1985) "Carbon Fibres and Their

Composites", ed. by E. Fitzer, Springer-Verlag, Berlin, p. 6Z 3- F. Tuinstra and J. L. Koenig (1970), J.Comp.Mat., voU p.49Z 4- C. Galiotis, R. J. Young, P. H. J. Yeung and Batchelder D.N. (1984):J.of

Mat.Sci., vo1.l9, pp.3640-3648. 5- C. Galiotis, I. M. Robinson, Young R.J., B. J. E. Smith and D.N.

Batchelder (1985): Pol. Comm. vol. Z6, pp. 354-355. 6- C. Galiotis and D. N. Batchelder (1988), J.of Mat.Sci.Let.,vol.7 p.545 7- R.J. Day, I.M. Robinson, M. Zakikhani and R. J. Young (1987), Polymer,

vol.Z8, p.1833 8- M.G. Dobb (1985): "Handbook of Composites", vo1.1,ch.17,p.673, Elsevier 9- H. L. Cox (1952): Brit. J. App1. Phys.,vo1. 3, p.72. 10- J. M. Whitney and L.T. Drzal (1987), Toughened Comp.,ASTMSTP937,p.179 11- C. Galiotis at al., to be published

769

Table 1. Relative intensities, frequencies and strain dependences for the prominent first order Raman bands of high modulus (HM) carbon fibres (Courtaulds). OX, IX signify oxidatively untreated and treated fibres respectively.

Fibre Mode

OX

IX

Normalised Intensity

0.20

1.00

0.40

1.00

1460

Raman Freq. (cm-1 )

1357

1574

1361

1578

1540 1620 1100 RAMAN SHIFT I em - , J

Strain (cm- 1/ %)

7.3'!: 0.6

9.1 1: 0.6

7.4 1: 0.7

8.71: 0.7

Fig.l-Raman spectrum in the first order region of a HM carbon fibre.

1585

1580

I E 1575 u

~ 1570

• load o unload

I I 1

FRACTURE

T

~1360~11 ~. A'9 Mode 1

~ 1355 •

1350 I

o 0·2 0·4 0·6 0·8 1·0

STRA1N t'lol

Fig.2-Raman Frequency as a function of uniaxial strain for the E2 and Al bands of a HM carbon fibre. The solid line is a least-squares lit to uil experimental data.

770

le .. ;..:R::am::.:a:::n.:....:.:.fr.:::eq:!:u:::e::.:nc::,y~(c::.m'-'----'1'--) ______ -,

181"

1812

,.,0

HI08

1108

o 0

o

Cycle 1

S • 4.37 (a)

1~.L-_~_~_~ __ L-_~_~~ o 0.' 0 .• t2 2.'

Strain (%)

Fig.3-Raman frequency as a function of strain for the 1615 cm- 1 band of a Kevlar 49 fibre.

"0

)( llS Eo.'

Q)

?-

F'lcl iun (Ktlly_ •• )

I ~:.-.... : :'Oj~d"'-• 0 I 0

00 00

0 0 I (><>0 I

o °0

~ ~

. .

I I I 0 00

0.0.". ~-:':-~:'-:-~,y.."'-::"':-'--7'c,--..-"L,----'.,-,J-.. ::';O"'-:-l

O~I/~lc~ Fibre length (mm)

Fig.4-Normalised axial tensile strain as 'a function of position along fibre length for a Kevlar 49/ Epoxy (Ciba Geigy 1927) system. The sample has been recycled three times.

fibre strain (,,) fibre slrain (%)

matrix strain elm)· 1.0% matrix strain elm)' 1.70 '10 ... . .. Ca) (b) ... . ..

... •.• !

.U liiM •..

o O..J: G.A 0.1 0 .. ,\2 \.Ao u o.A U Cl.t

fragment length (mm) fragment length (mm)

Fig.5-Fragmentation tests on a HMS4 (Hercules)/ Epoxy (Ciba Geigy 1927) system:

(a) fibre strain versus fragment length for 1% matrix strain (b) fibre strain versus fragment length for 1.7% matrix strain.

DEVELOPPEMENT DE METHODES DE CONTROLE PAR EMISSION ACOUSTIQUE DES STRUCTURES COMPOSITES

ABSTRACT

C. HERVE, M. CHERFAOUI, M. TRUCHON', X. DUFOUR*'

CETIM - 52 avenue Felix Louat - 60300 SENLIS - France "Elf Aquitaine - Boussens - 31360 ST MAR TORY - France

""SBPI - BP1 - 89220 ROGNY - France

Acoustic emission (EA) is a technique by means of which the inspection specifications (CARP Code) for composite materials have been produced over the last few years. This article describes the methodology used for implementing the "CARP" procedure concerning the qualification and in-service tests of tubular structures of glass fiber reinforced resin. The results obtained are also shown by comparing products from different suppliers.

1. INTRODUCTION

L'utilisation des materiaux composites dans plusieurs domaines (genie chimique, mecanique •.. ), la diversite des produits de base proposes, ainsi que l'absence de normalisation dans Ie CND rendent necessaire Ie developpement de methodes de contrOle non destructif fiables.

L'emission acoustique, qui designe l'onde de contrainte produite par la liberation d'energie au sein d'un materiau soumis a contrainte, presente dans certains cas les qualites necessaires pour constituer une telle methode. En effet, il existe aujourd'hui plusieurs normes et specifications pour Ie contrOle par emission des structures composites /1/.

772

Le present article presente les resultats des mesures d'emission acoustique effectuees sur les tubes lors de la mise en oeuvre de la procedure CARP (Comite d'Emission Acoustique dans les plastiques renforces) /2/. Les performances de tubes provenant de deux constructeurs differents (A et B) sont comparees et les criteres de qualification sont discutes.

2. I.E BANC D' ESSAIS HYDRAULIQUE

2.1. Caract~ristiques des tubes' tester

Deux types de tubes (type A et B) ont ete testes et sui vis par Emission Acoustique.

Ce sont des tubes de canalisation constitues de fibres de verre impregnees de resine epoxy. Leur pression de service est de 20 bars. Le diametre est de 150 mm, l'epaisseur mesuree de la paroi est de 4,6 mm (tube A) et 5,8 mm (tube B). La longueur des tubes testes est de 1 m.

2.2. La plate-forme hydraulique

Cette plate-forme garantit des vitesses constantes et reproductibles de pressurisation et de depressurisation. On sait, en effet, que la vitesse de sollicitation est un facteur d'influence essentielle pour l'emission acoustique. II est done necessaire d'en garantir la reproductibilite si l'on veut comparer les resultats d'essais successifs.

Pour effectuer les mises sous pression, on utilise une alimentation programmable qui commande la servovalve de la plate-forme hydraulique et qui est programmee par un calculateur.

Le cycle de.la mise sous pression utilise est celui du code CARP OU l'on monte jusqu'a la pression de qualification en pass ant par des paliers et des descentes en pression. Dans notre cas, la pression de service (Ps) est de 20 bars et la pression de qualification est fixee a 1,5 Ps (30 bars).

2.3. Surveillance par Emission Acoustique

Quatre capteurs resonnants a 180 KHz ont ete disposes sur Ie tube. Les parametres d'Emission Acoustique sont enregistres par un systeme multivoies de caracterisation et de localisation des sources d'Emission Acoustique. Toutes les voies d'acquisition sont calibrees selon la procedure CARP [3].

773

3. RESULTATS DES ESSAIS DE QUALIFICATION SELON LE CODE CARP

3.1. Definition des criteres du code CARP

Les essais de qualification des tubes ont ete effectues en utilisant Ie cycle de mise sous pression propose par Ie code CARP [3].

Au total 9 essais ont ete realises

3 essais sur tubes A sains

3 essais sur tubes B sains

A4, A5, A6

B3, B4, B5

3 essais sur tubes B ayant subi au centre un choc d'energie determinee (11 joules) : BD1, BD2, BD3.

Les criteres d'acceptation de l'essai de sont indiques au Tableau 1. II parait definition et la determination definie reprises dans la norme ASTM E 1119 [1].

qualification du code CARP important d'en rappeler la

dans la procedure CARP et

L'amplitude Am qui apparait dans Ie premier critere est la moyenne entre Ie seuil de detection et Ie seuil de reference de l'amplitude definie par la calibration de la chaine, dans notre cas Am = 60 dB.

Le deuxieme critere de comptage du nombre d'alternances Nc est determine en utilisant la rupture d'une mine de crayon en deux points de calibration. La distance entre la source et Ie point de calibration sera choisie de sorte que l'amplitude des signaux recueillis soit egale a l'amplitude Am definie precedemment. La valeur Nc a ete determinee pour chaque tube et portee dans Ie tableau 2.

Le troisieme critere concerne l'evaluation de la duree des salves enregistrees qui ne doit pas depasser la valeur M definie par calibration. Cette valeur correspond a des signaux d'amplitude moyenne et de longue duree qui sont associes au delaminage ou a la rupture de joints colles. Ce critere ne sera pas pris en compte car nous n'avons teste que des tubes droits.

Le quatrieme critere concerne les salves dont l'amplitude depasse Ie seuil d'amplitude de reference, dans notre cas 72 dB. Ce critere est souvent associe a la rupture de fibres et donc significatif de dommages structuraux majeurs.

Le cinquieme critere est fonde sur l'evaluation du rapport "Felicity" qui est defini comme Ie rapport de la charge de reprise de l'emission sur la charge maximale atteinte precedemment. Ce critere n'est pas applicable lors des premiers chargements.

774

3.2. Interp~tation et comparaison

Dans Ie Tableau 2 sont notees, pour chaque essai, les valeurs des parametres d'Emission Acoustique correspondant a chaque critere.

a) Les tubes B

On cons tate que les trois tubes B sains testes ne repondent pas aux criteres de qualification du code CARP. C'est toujours Ie premier critere qui n'est pas verifie il y a emission de salves dont l'amplitude depasse 60 dB apres les quatre premieres minutes de palier. On notera que ces salves apparaissent, sauf pour B3, au palier de 30 bars qui est la valeur de la pression maximum de ce cycle. Tous les autres criteres sont verifies.

Pour les tubes B, ayant subi un choc, les valeurs obtenues sont nettement differentes de celles des tubes sains. Leur emissivite est plus precoce et plus importante, comme en temoignent les valeurs des deux premiers criteres. En plus, Ie critere 4 n'est pas toujours veri fie , sauf pour Ie tube BD2 qui, apres analyse, a montre qu'il portait une couche supplementaire.

Le code CARP permet donc de bien distinguer des tubes sains et des tubes de m~me type comport ant des defauts.

b) Les tubes A

lIs verifient tous les criteres du code CARP. Les trois tubes testes sont donc qualifies. On remarque m~me que Ie tube A6 n'a donne lieu a aucune Emission Acoustique durant Ie cycle de chargement.

3.3. Discussion des crit~res du code CARP

L'application des criteres de la procedure CARP permet bie~ de separer les tubes en 2 populations : les tubes B non qualifies et les tubes A qualifies. Cependant, ces criteres n'interviennent pas tous dans la decision (Tableau 2).

Le critere 1 est Ie plus discriminant pour separer les et B; c'est lui qui a permis dans la plupart des cas decision. Ce critere est associe a la micro fissuration II est donc probable que la mat rice soit 1 'element difference entre les tubes.

tubes sains A de prendre la de la matrice. qui fasse la

Le critere 2 presente une grande dispersion qui est a rapprocher des problemes de propagation d'onde. Effectivement, Ie nombre d'alternances depassant un seuil fixe est fonction de l'attenuation du signal recueilli. Ce deuxieme critere doit ~tre mieux defini pour ~tre plus exploitable.

Le critere 4 concernant les salves d'amplitude superieure a 72 dB est souvent associe aux ruptures de fibres. II prend son importance notamment pour les tubes avec defaut. II parait comme etant Ie critere Ie plus severe.

775

4. CONCLUSION GENERALE

Des essais de mise sous pression ont ete realises sur des tubes droits en verre-epoxy utilises dans l'industrie petroliere. Le SU1V1 par Emission Acoustique a permis de tirer plusieurs conclusions :

1 - Concernant les essais de qualification selon la procedure CARP, tousles tubes B ne sont pas qualifies pour une pression d'epreuve de 30 bars, contrairement aux tubes A, alors que ces tubes sont consideres equivalents par leurs constructeurs.

2 - Les tubes B ayant subi un choc ne verifient pas egalement lescriteres du code CARP ils ne sont donc pas qualifies pour une pression de 30 bars. L'emission est plus importante que dans les tubes sains de m~me type.

3 - L'absence d'une couche est bien mise en evidence par l'Emission Acoustique.

4 - Des essais futurs doivent permettre d'etudier l'influence des assemblages colles et une meilleure definition des criteres et de la calibration du code CARP.

BIBLIOGRAPHIE

[1] M. CHERFAOUI, A. LEMASCON et J. ROGET ContrOle par Emission Acoustique renforcees : reservoirs, capacites et JEC 1987 - Composites N° 3 - Mai/Juin

[2] M. DROGE

des structures tubes. 1987.

plastiques

Recommended practice for Acoustic Emission testing of fiber-glass reinforced plastic piping systems 1st Int. Symp. on Acoustic Emission from Reinforced Composites. SPI - SAN FRANCISCO. July 1983. Session 4 (voir aussi traduction CETIM).

[3] M. CHERFAOUI, C. HERVE, J. FEBVAY, X. DUFOUR Developpement de methodes de contrOle des structures composites JEC 1988 - Composites N° 3 - Mai/Juin 1988.

776

Tableau 1 : CRITERES P'ACCEPTATION [21

Une structure RTRP acceptable devra veri tier tous les criteres ci-des80us

Cri teres Tubes neufa ce et/ou structures SIGNIFICATION

tu a 34 MPa PU CRITERE

--1 PREMIER CHARGEMENT CHARGEMENTS SUIVANTS

Emission en Pas d' ampli tude Am et - Pas d'emission apres Mesure de l'endommage-

palier TH mn de palier lIent permanent dO i la

- pas plus de 5 evene- lIicrotissuration de la

menta par minute .atrice

apres TH lIinutes

2

Nombre total - interieur a Nc - interieur i Nc/2 Mesure l'endommagement

de coups global pendant Ie cycle de charge

--3

DBC - inferieur a M - infiorieur a M Mesure le delaminage,

la rupture de joints

i la croissance de

fissure importante

-- --4

Nombre - interieur i 10 - interieur i 10 Mesure des ruptures

d'ivenements microstructurale. de

dont l' ampli- forte energie

tude ) Ce critere est souvent

ampli tude de aasaeie aux ruptures

reference de tibres

5

Rapport - superieur i 0,95 - superieur i 0,95 Mesure 1a gravite des

"Felicity" <si applicable) endommagements intro-

duits anterieurellent

'----

777

Tableau 2 Critares du code CARP pour les preaiers charlementa

Critere I Critere 2 Critere 4 Decision

Type Nombre de salves a Nombre total de Nc Hombre No.bre de

de Am = 60 dB apres salves apr .. 2 mn de total salves

tube 2 mn de palier et palier (A l 50 dB) d' alter- l 72 dB

valeur du palier

en bars

1---

B4 I salve Ii 30 bars 14 3810 2140 8 Non qualifie

B5 2 salves Ii 30 bars 6 2485 342 0 Non qualifie

BDl 5 a 22.5 bars 109 8130 5176 9 Non qualifie

BD2 2 a 22.5 bars 21 7980 1029 6 Non qualifie

BD3 9 a 30 bars 85 10820 4814 10 Non qualifie

A4 0 0 43530 1082 0 Qualifie

A5 0 0 865 224 0 Qualifie

A6 0 0 16664 0 0 Qualifie

--- --

CURE CHARACTERISTIC DETERMINATION USING MICROELECTRONIC DIELECTRIC SENSORS

D. DAY, H. LEE, K. RUSSELL, J. WHITESIDE

Micromet Instruments Inc. University Park - 26 Landsdowne street, Suite 150 - 02139 CAMBRIDGE, MA -USA

The development of microdielectrometry has made possible production monitoring and cure control in a variety of thermoset resin based systems. As the demands of the technology grow, existing sensors must be demonstrated to have applicability or new sensors must be designed.

This paper describes recent microdielectric applications to unusual materials or processes. Among new applications of existing sensors are polyurethane foams, pultrusion forming, and adhesive composition control. A new ceramic sensor has been developed. It is operable at temperature in excess of 500°C and designed for repetitive use in a tool or press. The use of this sensor to evaluate both curing and cured materials is also reviewed.

1 - INTRODUCTION

Microdielectrometry was initially applied to well behaved thermosetting resin systems, typically epoxies. With such resin systems reaction temperatures were usually below 175°C and pressures were low to moderate. But as the technology of high performance composites developed, there have been growing requirements for extending the microdielectric measurement technique to a broader range of materials and processing conditions.

One of the major drivers for these requirements is that the microdielectric measurements can be made both in the laboratory and in the production environment with the same equipment. They form, therefore, a bridge between the sensitive and vital laboratory equipment and the heretofore limited process control instrumentation; the clock, thermometer, and pressure gauge.

The purpose of this paper is to review this extension of technique and illustrate results with typical data.

2-BACKGROUND

The dielectric properties of a material can be measured by monitoring its response to an oscillating electric field. Recent developments in the field of microelectronics now enable the fabrication of highly sensitive implantable microdielectric sensors (1,2). Not only are the sensors small (active area = 2 x 3 mm, see Figure 1) but they also function down to very low

780

frequencies (.005 Hz). The lower frequencies are important to monitor polymerization reactions near completion when relaxation times grow long.

The two dielectric properties are the permittivity (frequently but incorrectly called the dielectric constant) and the loss factor. They are the result of dipole motion and ionic conduction. In addition, the measurement of the dielectric properties can be influenced by electrode polarization.

The changes in permittivity observed during typical polymerization reactions are only a function of dipole motion. As reaction proceeds, the motion of dipoles is hindered and the permittivity falls. The usual change in permittivity is a factor of 3-6.

The dielectric loss factor is influenced by both dipole motion and ionic conduction. The dipole motion contribution is relatively small throughout typical polymerization processes. The ionic conduction contribution is the result of motion of ionic impurities in the material. As little as 1 part per million of such impurities will produce significant ionic conduction levels (3). Since ionic conduction is essentially a measure of the mobility of ions in the resin, it can be correlated with viscosity prior to gelation, and to degree of cure after gelation. It must be noted that gelation is not a dielectric event; it is merely a step on the continuum from uncured to cured material. The correlation is strong because ionic mobility, and, therefore conduction, usually changes by several orders of magnitude during the course of the polymerization process. Further, since the ionic conduction term is inversely proportional to frequency, it can always be made to dominate the loss factor if measurements are made at low enough frequencies. Loss factor is thus very sensitive to the degree of polymerization throughout the entire reaction. Implicit in this view is the assumption that the concentration of the ions does not change as the viscosity of the resin changes. Since ionic impurities generally do not enter into the polymerization reaction, this assumption is valid.

The phenomenon of electrode polarization arises from the interaction of ions in the test material with the electrodes used to apply the oscillating electric field. The maximum potential of the oscillating field is deliberately kept low to prevent electrochemical reaction. Therefore, ions build up at the electrode surface because of the applied field, but they do not discharge. The resulting charged layer acts as a large capacitor. This causes the measured loss factor, primarily induced by ionic conduction, to be lower than otherwise and, at the same time, the apparent permittivity is abnormally elevated. A detailed discussion of this effect is given in Reference 4.

Figure 2 shows typical dielectric data for a commercial polyurethane repair adhesive. The influences of dipole motion, ionic conductivity, and electrode polarization on loss factor are shows at several frequencies. Initially, loss factors are high due to high ion mobilities in the low molecular weight resin. As the resin reacts the loss factors decrease and each frequency in tum (from high to low) exhibits a dipole loss peak.

3 - DIELECTRIC MEASUREMENTS

The measurements reported here were made with the Eumetric System II Microdielectrometer manufactured by Micromet Instruments, Inc. Unless otherwise noted, disposable, integrated circuit sensors featuring built-in thermal diode for temperature recording, were used to monitor all dielectric responses. The instrument is capable of simultaneous and independent measurements of both permittivity and loss factor at frequencies ranging from .005 to 10,000 Hz. The instrument is driven by an IBM PC or equivalent using software provided by the manufacturer.

4 - DISCUSSION

4.1. Correlation with Viscosity The relationship between dielectric loss factor and viscosity is so important for process

control considerations that a brief review of supporting data must be included in any discussion of dielectric measurements.

Figure 3 shows the relation between dielectric loss factor and viscosity for polystyrene, a simple thermoplastic. The viscosity data were calculated from temperature using an empirical relationship provided by the resin manufacturer, General ElectriC. The 10 Hz loss

factor data were conductivity dominated so that the ordinate, labeled "Loss Factor" is equivalent to conductivity. The reciprocal relationship between the viscosity and the loss factor, which is conductivity dominated, is expected.

781

This correlation also appears to hold in the case of reacting thermosetting resins. A Rheometries Dynamic Spectrometer (RDS) was used by the McDonnell Aircraft Coll1)any to measure viscosity of samples of neat epoxy resin (Hercules 3501-6). SalJ1)les from the same batch of resin had been evaluated about five months earlier using the micI'odielectrometry technique (5). The viscosity data are presented in Figure 4 as fluidity to permit more ready COll1)arison with the dielectric loss factor.

The agreement between the two techniques in terms of magnitude of change and rate of change is excellent.

4.2. Polyurethane Foam Dielectric data taken during the polymerization of soft polyurethane foam used in the

automotive industry. The reaction is fast and most of the physical (and dielectric) change occurs in the first two or three minutes. The reactions were carried out in an aluminum mold measuring approximately 30 x 30 x 2.5 cm (12" x 12" x 11 and preheated to 38°C. The prewelghed isocyanate-urethane components were mixed in a high speed blender and poured into the mold. Pouring was completed within 10 seconds of the time the blender was turned on. Zero time was taken to be the starting time of the blender. After pouring, the lid was clamped in place and the reactions permitted to run to completion.

Four different proprietary forrrulations, having differed cure speeds, were tested. The conductivity data were extracted from the loss factors using methods described by Day et al (4). The time derivative of the conductivity, which may be considered as the "Dielectric Reaction Rate; is shown in Figure 5 for the four formulations tested. Hs greatest change takes place during the first 2-3 minutes. Then, as expected, it approaches zero. The relative reaction rates determined in this manner were in complete agreement with the rates as determined by conventional methods. Similar variations in reaction rates were observed with a single formulation reacting at different temperatures.

Production efficiency requires that the time spent in the mold be just long enough to assure dimensional stability of the part after it is removed from the mold. Too much time spent in the mold reduces throughput; too little time resuHs in rejected parts. The necessary time varies with both forrrulation of the material and temperature of the molds and reactants. Formulations can be maintained relatively constant. However, the temperature of the molds and materials is determined almost completely by the armient conditions of the production floor. From the data shown above it is evident that the dielectric monitoring can provide the necessary cure time information on lne and in real time. An exa",,* of how data like this can be used for closed loop cure control has been described by Bromberg et al (6).

4.3. Pultrusion The puHrusion process required careful process control to insure integrity of the final

product. However, the common measurements of tefTl)erature and feed rate do not address the characteristics of the polymerizing material; in fact, the assumption is made that the reaction mixture in invariant. AHhough this assu!Tl)tion is not valid, it may be satisfactory as long as the final properties of the polymerized material are tolerant of process parameters. As specifications become more restrictive, it is expected tbat closer control of the process will be required to achieve satisfactory end product.

The applicability of the implantable Eumetric sensor to pultrusion was demonstrated with a polyester-glasscomposite. The 38 cm (15 in) long sensor was soldered to a 300 cm (10 It) ribbon cable which, in tum, was connected to the microdielectrometer. The sensor was then fed through the puHrusion die. It monitored the course of the reaction through completion until it reached the end of its travel at which point the ribbon cable was cut. The sensor and cable were disposed of along with a length of the extruded form into which they were embedded.

Data for one of these runs are shows in Figure 6. The data were processed in terms of degree of cure and viscosity. Also shows is the temperature recorded by the integral thermal diode on the sensor. It is evident that this technique offers a method for control of the

782

process by using predetermined parameters to control the throughput speed and/or the temperature of the die. Specifically designed sensors would be permanently mounted in the die to monitor the process and control it though appropriate actuators. The closed loop control techniques of Reference (6) could be applied to such a process.

4.4. Composition Monitoring Day has shows that dielectric properties can be used to verify the composition of

reacting mixtures of epoxy resin and amine hardener by comparing the dielectric loss factors near the end of cure (6). There is interest in showing if composition can be determined early in the reaction; particularly, at the point of mixing. The interest arises because on-line mixers generally use speed control to obtain correct mix ratios. However, wear of pumping and mixing elements can cause deviation from desired composition limits. A measurement of the mixture properties that can be related to composition would provide a closed loop control mode that would assure that correct mixture rations are met despite wear and age induced changes in critical feed components.

Mixtures were prepared of a commercially available 2-part adhesive in which the ratio of the two parts was varied around the recommended mix ratio. The mixtures were prepared by preweighing the two components, mixing them rapidly by hand, and then placing a sample of the mix on the active surface of a microdielectric sensor. The data are shown in Figure 7. Despite the difficulty involved in uniformly mixing two rather viscous materials and in making the measurement before significant reaction has taken place, the dielectric response appears to be linear over the range investigated and may be large enough to provide a feedback control signal.

4.5. Ceramic Sensor The disposable integrated circuit microdielectric sensor is limited to use at about 300°C.

In addition, there are applications in which the use of this sensor for reasons such as pressure, part of mold configuration, or cost, is not appropriate. Other sensors ITlJst then be used. One of the recent developments is a ceramic sensor that can be used at temperatures in excess of 500°C and pressure limited only by its support structure.

The ceramic sensor active element is approximately 2.5 x 2.5 em (1" x 1") and it Is in the form of a metallized interdigitated array with line width and spacing of 0.25 mm (.010 in). The ceramic base is made of high density alumina about 0.25 mm (.010 in) thick. For experimental purposes the sensor is supported on a stainless steel platform in which is provision for temperature instrumentation.

In order to test the performance of the ceramic sensor, a cure was conducted using both the ceramic sensor and the disposable integrated circuit sensor. The material that was cured was Fiberite 8-212 graphite/epoxy prepreg. The cure schedule was a ramp to 175°C (350°F) at 2.2°C/min, a hold for 2 hours, and cool down to room temperature. Temperature was measured using the temperature sensor in the integrated circuit sensor. The resuHs of the two simuHaneous tests are shown in Figure 8. The two curves reach their conductivity maxima (viscosity minimum of the resin) at the same time. In addition, each curve reaches a slope of essentially zero (end of reaction) at the same time. The two sensors are therefore providing equivalent information despite a sensing area ratio of more than 200. The analytical procedure used to convert the basic gain and phase measurements into loss factor and permittivity is thus well validated.

At the conclusion of the cure described above, the integrated circuit sensor was completely bonded into the composite. The ceramic sensor, whose surface had been treated with a conventional silicone based mold release agent, was easily separated from the cured composite and immediately ready for another use.

5-SUMMARY

Microdielectrometry has application to a wide variety of thermoplastic and thermosetting resin systems and processes. Measurements can be made in high speed reacting systems, such as polyurethanes, or in completely static systems, such as cured solids. Composition variations can be detected at the beginning of cure. Dielectric data are shown to be correlated

with viscosity and with degree of cure. This capability is can be applied to monitoring and control of both conventional batch thermoset cures as well as continuous pultrusion. A reusable ceramic sensor, which is equivalent in measurement capability to the integrated circuit miniature sensor, has been demonstrated.

REFERENCES

783

1. S.D. Senturia, N.F. Sheppard, H.L. Lee, and D.A. Day, "In Situ Measurement of the Properties of Curing Systems with Microdielectrometry," Jrnl. Adhesion, 15,69 (1982)

2. N.F. Sheppard, M. Coin, and S. Senturia, "A Dielectric Study of The Time Temperature Transformation (Tn) Diagram of DEGBA Epoxy Resins Cured with DDS," Proc. 29th SAMPE Symposium, Reno, NV, 1243 (1984)

3. A.A. Blythe, "Electrical Properties of Polymers," Cambridge University Press, Cambridge.

4. D.R. Day, T.J. Lewis, H.L. Lee, and S.D. Senturia, "The Role of Boundary Layer Capacitance at Blocking Electrodes in the Interpretation of Dielectric Cure Data in Adhesives," Jrnl. Adhesion, 18, 73 (1984)

5. "Computer Aided Curing of Composites," 5th Quarter1y Interim Technical Report, Contract F33615-83-C-5088, McDonnell Aircraft Company, SI. Louis, MO, July 1985

6. M.L. Bromberg, D.R. Day, and K.A. Snable, "Measurement and Application of Dielectric Properties," Electrical Insulation, 2,18 (1986)

7. D.A. Day, "Effects of Stoichiometric Mixing Ratio on Epoxy Cure: A Dielectric Analysis," SPE Technical Papers, 31, 327 (1985)

:.

~~g~~11tr .1 10 JI1 100

-::::::===========± lK 10K . 010" -1 L-_--'-__ -'-__ "--_---' __ -'-_--.J

-Figure I . Inlegraled ClrcuH Sensor

o 50 100 150 200 250 300 Tlmo(mlnl

Figure 2. Polvurelhene Repelr Resin ,2JO"C]

Slo~. -1 ,04

/ C4 ... COlIn :; ,"2

784

.. ·f I ,~ ~

":r I .'::" ,. .. r I~-"I

'­- .

.. r .. IL---_-'--_"'""'--__ .....::::= ....

Figure 4. Viscosity - Loss Factor Correlation For Epoxy

I . I J 1

' ,1 ---==~;;;;~I I

i I !

Figure 5. Reaction Rates of 4 Polyurethane Foam Formulations

.. o o 0

...

.. '--______ "-_....".. ___ ----J

u.o 40.0 EpoE., U.CI "-u..o .10.0 ~ IU"

Figure 7. Dielectric Me.,uremen ... of Mix Ratios

100 l 200 ..",.------., /' .....

/ ..... I

10 ,OO T_p. ;'

e / !l e / r 10 120 u / • ~ u / t- 40 '0 /

/ Ion VllCoalty .' . ". / . ' . / ....

20 "2--- /

•

Figure 6. Pultrusion Monitoring

J - .. f , I :r .' [! '4\,L ~.::::--\==' --.-c-...w-

_11. '" - .,-.. -----; .. ;--=----::::----:, ..

Figure 8. Sensor Comparison (Fiberite 8·212)

• .i ;;

3 , <

~ z ~

ABSTRACT

THE USE OF LASER MOIRE INTERFEROMETRY IN THE STUDY OF DEFORMATION FIELDS

IN COMPOSITES AND ADHESIVES

R. DAVIDSON

HafWell Laboratory - B 47 Materials Development Division HafWell Oxon - OX11 ORA ABINGDON - England

Laser moire interferometry is a recently developed optical technique which allows high resolution measurements of deformation fields in stressed materials. This paper describes the technique and illustrates its use in strain concentration measurements around circular holes in CFRP laminates and in a stressed adhesive bond. In the latter example the initiation and growth of damage have been successfully followed.

1. INTRODUCTION

A high resolution laser moire technique is being used at Harwell in the study of strain fields in materials. This technique is particularly applicable to composite materials where rapidly varying strain fields are often present. In these materials it is essential to take into account the effects of secondary stresses in both materials and component design. These stresses can lead to premature failure of components and are usually related to out of plane, transverse or shear loading. They can be affected by the laminate stacking sequence, cure schedule, matrix material and hybridization. Joining composite materials by mechanical or adhesive bonds and the effects of holes and notches gives rise to stress concentrations which must also be taken into account in design.

2. LASER MOIRE INTERFEROMETRY

Moire techniques rely on a parallel grid, carried in the surface of a specimen, which deforms when the state of strain within the

786

specimen is changed through, for example, mechanical or thermal stress. Interrogation of the specimen grid through an undeformed aligned reference grid produces contour maps of equal displacement -moire fringes (1). Until recently the technique was limited by the coarseness of the finest grids achievable by ruling methods (max. 40 lines/mm). Post (2) first succeeded in identifying a suitable system to produce grids with frequencies up to 2000 lines/mm using a photographic recordJng of two beam interference of coherent laser light. A series of parallel lines of light and dark fringes are produced over the exposed area. On developing, the relative shrinkage of the unexposed and exposed parts of the photographic plates produces a corrugated phase type grating. Cross gratings can be produced by rotating the photographic plates through 90° between a double exposure.

The gratings are replicated onto the surface of the specimen to be examined. This is achieved by evaporating a thin aluminium layer on to the master grating which is transferred to the specimen via a low modulus photoelastic epoxide curable at room temperature. When the adhesive has cured the master grating is easily prized off to leave the replicated aluminium layer adhering to the specimen surface. The success of the technique relies on the production of a virtual reference grating being generated in space which interacts with the deformed specimen grating to produce deformation fringes. Orthogonal reference gratings are made using appropriate interference of expanded laser beams. Two expanded beams of coherent HeNe laser light crossing at an angle of 2 ~ in a horizontal plane form a virtual reference grating with vertical lines. In the present work using laser a wavelength of 633 nm and a grating frequency of 1200 lines per mm ~ is set at 49.4°. Similar intersecting beams in a vertical plane gives horizontal lines. These conditions are achieved using a parallel laser beam set at angle ~ to the specimen expanded so as to illuminate the specimen and one plane mirror orthogonal to the specimen axis and two plane mirrors at ± 45° to the axis. Each reference grating interacts with the similarly orientated lines of the specimen grating to form moire patterns observed when imaged orthogonal to the specimen surface. The patterns give maps of constant displacement

u = 1 N and v = 1 N f x f y

where u and v are the in-plane x and y components of displacement at any point on the surface and Nx and N are the corresponding fringe orders. f is twice the specimen grat{ng frequency and is equivalent to 2400 lines per mm. Strain information can be obtained by applying Lagrangian relationships

au + 1 «au)2 + (av)2) 1 aNx EX = '" - (-)

ax 2 ax ax f ax

av + 1 ((au)2 + (au)2)) 1~ Ey = '" - ( ) ay 2 ay ay f ay

(~ + ¥X) 1 aNx ~ "6 xy = arc sin '" - (- + ) (1 + EX) (1 + Ey) f ay ax

The approximations at small strains are also given EX and E are normal strains in the x and y directions respectively and "6~y is shear strain.

787

A specially designed loading frame has been constructed to be used in conjunction with the required optics. Accurate mirror positioning is achieved by mounting on small goniometer stages. The loading frame is seated on a tilt and rotation stage, and the load is applied through a worm drive mechanism and is reacted by two aluminium alloy columns. The load frame and associated optics are mounted on standard triangular optical bench rails. A 7 mW HeNe laser is diverged through a 60x microscope objective and a spatial filter at the focus is used to clean the diverging laser beam, which is collimated with a 150 mm diameter planovex lens 'of focal length 380 mm. Deformation fringes were recorded photographically using 9 x 11 mm Polaroid film. Alternatively video techniques have recently been applied. This is particularly appropriate for real time measurements of creep and damage development.

3. EXAMPLES OF THE USE OF LASER MOIRE INTERFEROMETRY

3.1 Strain concentration effects in composite laminates

Figures 1-3 show the u and v displacement fringes around 3.1 mm diameter holes in 25 mm wide XAS CFRP laminates constructed of unidirectional reinforcement, (0/90)s and (0/90) 5 harness satin woven fabric material respectively. The unidirectional material is clearly more anisotropic and very large fringe gradients are maintained around the hole compared to the cross-ply laminate. The results in Fig. 3 indicate how the woven structure of the reinforcement has a marked effect on the localised deformation fields. The tensile concentration factors estimated from the moire data are given in Table 1, along with the predicted values using the theories of linear anisotropic elasticity discussed by Lekhnitskii (3). The stress concentration factors tangental to the hole are predicted to be

The measured maximum stress concentration factor in the unidirectional material agrees to within 7% of the predicted value. The measured value for the cross-ply material is 30% lower whilst the measured value for the woven material is 20% in excess of the predicted value. These differences are significant and are being investigated further. Recent work by Post (4) has also identified

788

some unusual effects, in particular, the strain concentration effects in certain laminates have been observed to vary significantly with stress.

3.2 Applications to the study of structural adhesives

Figure 4 shows the v displacement fringes in a 0.53 mm thick epoxide adhesive, in a thick adherend shear test, at load levels of 3 kN, 3.75 kN and 4.75 kN. The specimen geometry is also shown. By introducing a carrier pattern of rotation the fringe density along the loading axis can be changed to aid interpretation. Average peel strains across the thickness of the adhesive, along the bond length can be calculated from the fringe displacements,~. The peel strains are compressive in the central region of the bond length and become tensile as the cut outs are approached, with the maximum occurring at the cut outs. At 3 kN the fringes are continuous over the whole bonded region. The peel strain concentration at the cut out is clearly seen in Fig. 4b. At loads greater than 4 kN small interfacial adhesive cracks are observed growing from the cut outs. The cracks are stable and extend only with increasing load. Plastic deformation is also observed in the steel adherend adjacent to the cut outs as in Fig. 4c.

The shear strains across the adhesive can be obtained from the ~ values obtained from the data as shown in Fig. 5. The fringe gradient is greater close to the cut outs. Quantitative measurements of shear and peel strains measured using laser moire have been used at Harwell to refine computer programs that predict the strain distributions and strengths of structural adhesive joints.

4. CONCLUSIONS

• Laser moire interferometry has been shown to be a unique and invaluable technique for the study of full field deformation in composite materials. The technique offers the only quantitative method available for the full-field study of secondary stress effects in composites.

• The technique is invaluable in the quantitative measurement of highly localised strains in adhesive bonds and aids the development of theoretical predictive models for joint strengths. The initiation and progression of damage has been followed.

5. ACKNOWLEDGEMENTS

This work was undertaken as part of the UKAEA underlying research programme.

789

6. REFERENCES

1 - Thearcaris, P.S., Moire Fringes in Strain Analysis, Pergamon Press, 1969.

2 - Post, D., Moire Interferometry at VPI and SU Experimental Mechanics, Vol. 23, No.2, June 1983, p. 203-210.

3 - Lekhnitskii, S.G., Anisotropic Plates, Ed. S.W. Tsai and T. Cheron, Gordon and Breach, Science Pub] ishers Inc., New York, 1968.

4 - Post, D., The Analysis of Deformations and Strains in Composites by Moire Interferometry, ICCM VI and ECCM 2, Edited by F.L. Matthews, Vol. 5, p. 251-261, 1987.

Table 1 Data for CFRP laminates

X-AS El E2 G12 "12 k90 kO k90 kO Composite

Lay-up GPa GPa GPa Predicted Measured

0° 127.6 6.0 5.5 0.249 6.65 -0.22 6.25 -0.6

(0/90)s 67.0 67.0 5.5 0.022 4.76 -1.0 3.3 -1.0

(0/90)woven 67.0 67.0 5.5 0.022 4.76 -1.0 6.0 -1.0

790

Figure 1 v and u fringes in a unidirectional CFRP laminate

loaded to 3.8 kN.

Figure 2 v and u fringes in a (0/90)s CFRP laminate

loaded to 2.5 kN.

Figure 3 v and u fringes in a woven CFRP laminate

loaded. to 3.5 kN.

a) 1000 N I

Figure 5 u-Moire displacement fringes at the central region

of the T AST joint as a function of load.

791

792

a) 3000N

CUTOUT

ADHESIVE

saUARE--"''''''''''' MOIRI:

GRATING

)(,u

Ly .. STEEL

ADHEREND

ADHESIVE TEST SPECIMEN

GEOMETRY

}

-~OIld I .... r---

b) 3750N C) 4750N

Figure 4 v-moire displacement fringes close to the cut out of a

T AST joint as a function of load.

NDE OF THICK GFRP COMPOSITES THROUGH ULTRASONIC WAVEFORM DETECTION

R. TETI, G. CAPRINO

University of Naples - Department of Materials and Production Engineering Piazzale Tecchio 80 - 80125 NAPOLI-Italy

ABSTRACf

Ultrasonic non destructive evaluation of thick glass fiber reinforced plastic composites was carried out using a complete waveform digital acquisition and analysis technique. It is shown that this method yields detailed information on the entire material volume, providing for the accurate localization of defects. The availability of complete waveform data files makes it possible to generate images acting as interfaces between user and ultrasonic database and to statistically process the data to obtain information on the material global properties. Material quality, determined through this technique, was found in good agreement with the results of macroscopic examinations.

INfRODUCflON

Non destructive evaluation (NDE) of composite materials through ultrasonic (UT) methods has been proved to be particularly effective in the identification of a variety of defects such as delaminations, cracks and localized or diffuse porosity [1 - 3]. It is generally thought that only relatively small thicknesses may be reliably tested using ultrasound because of the highly attenuating nature of composites. However, in [4] it was shown that, even in the case of glass fiber reinforced plastic (GFRP) composites, considerable information on the material structure could be obtained for thicknesses up to 50 mm through pulse-echo UT techniques; the use of digital UT techniques and advanced image processing [5] could significantly enhance the resolution of the NDE system.

794

The technique utilized in [5], based on image formation through single gate setting on the UT waveform [3], does not allow for a three-dimensional evaluation of material quality, unless several successive time-consuming scans are carried out to cover the entire material thickness using different gates. To obtain a volumetric evaluation of the material through only one scan, a technique of UT digital analysis based on complete waveform acquisition [6] may be used. In this case, a volumetric UT data file is formed, from which information on different thickness portions of the material under examination can be obtained via software.

In this work, an UT NDE technique based on complete waveform digital acquisition and analysis is applied to thick GFRP composites in order to determine its advantages and limitations.

I - MATERIALS AND EXPERIMENTAL PROCEDURES

E-glass mat/polyester resin laminates were fabricated by hand lay-up, followed by hot pressing. Four 8 mm thick samples of different quality and in-plane size 120 x 120 mm2 were cut for UT evaluation. The samples were subjected to macroscopic examinations to identify possible defects. Accordingly, a judgement on sample quality and structural integrity is reported in Tab. 1.

Pulse-echo ultrasonic C-scans were carried out using a 15 MHz focused immersion transducer. Scan area was 50 x 60 mm2 and scanning step 0.35 mm. During the scan, the entire RF UT waveform was digitized at 100 MHz sampling rate to fulfill the Nyquist criteria. The digitized waveforms were sent to a 3-D advanced graphic computer for data acquisition, storage and analysis. This procedure created very large UT data files ( .. 8 Mb). To generate images, a reference (mode) waveform was created by statistical processing of the· UT data. A time gate corresponding to a selected thickness portion was set from the keyboard on the digitized waveform. After gate setting, the deviation of the gated portion of each waveform in the file from the reference waveform was evaluated [6]. This value was utilized for image formation using standard software algorithm for conventional single-gate UT evaluations [7]. Different gates were set via software using the same UT database, without the necessity for multiple time-consuming C­scans, in order to examine the entire laminate thickness. Images were presented on screen and photographed through a hardcopier.

795

II - RESULTS AND DISCUSSION

In Fig. 1, images obtained from sample R7 are reported. To generate the four images, a typical waveform, retrieved from the UT data file, was presented on screen in a grid and four successive gates were set via software to subdivide the sample thickness into four equal parts. The referencing algorithm detailed in the previous section was then applied to each gate to create the corresponding image. Each image refers to a thickness portion of '" 2 mm.

The images represent active interfaces between user and UT data file [8]. Waveforms corresponding to any point within the image may be retrieved from the UT database by simply pointing and clicking the mouse at the relevant position. This allows for real-time acquisition of information about the entire material structure.

Fig. 2 reports the waveforms corresponding to the arrow locations in the images of Fig. 1. Fig. 2a represents the UT waveform corresponding to arrow (a) in Fig. 1. The position of the first gate is shown by two vertical bars in the grid. The waveform indicates a defectless material zone: internal peaks are not observed and the back echo is clearly visible. Fig. 2b shows the position of the second gate and the UT waveform detected at the location of a porosity. In this case, a peak is visible within the gate: its value determines the dark spot pointed by arrow (b) in Fig. 1. Peak position indicates that the porosity is localized approximately at the laminate midplane. The back echo is totally absent as the UT signal energy is completely reflected by the defect. Fig. 2c displays the third gate and the UT waveform detected at arrow (c) in Fig. 1. Macroscopic examinations had shown in this material zone an extended delamination (Tab. 1). Also in this case, the defect is signalled in the waveform by a high amplitude internal peak and the absence of back echo; in the image, it is indicated by a dark area. Fig. 2d shows the fourth gate and the UT waveform detected at a sub-surface porosity (arrow (d) in Fig. 1). The lower right image in Fig. 1, corresponding to this gate, displays an evident dark area of extended size designating the delaminated zone. The delamination determines the absence of back echo in the fourth gate, producing a high deviation from the reference waveform that yields a dark area in the image. Good quality material is characterized by a low deviation, yielding light grey areas.

The examination of the four images from sample R7 allows for some initial considerations. By comparing the UT waveforms in Figs. 2b and 2c, it may be noted that the analysis of a single waveform is not sufficient to discriminate between porosities and delaminations.

796

In both cases, the defect is "seen" by the UT waveform as a lack of structural continuity. Only the examination of the entire UT image allows the indentification' of the defect nature.

Besides the identification of defect nature and in-plane location, information about defect position within the thickness may be obtained. The image generated through the fourth gate (lower right in Fig. I), for example, indicates the presence and contour of a delamination; to identify its position within the thickness, a more detailed examination is necessary. The two upper images in Fig. I (first two portions of material thickness) do not evidence any lack of continuity at their lower right corner, where the delamination is. The lower left image in Fig. I shows concentrated dark spots in the delamination zone. The lower right image in Fig. I is formed through back echo gating and accounts for the attenuation induced by all defects encountered by the UT signals before hitting the back surface. Therefore, the defect may be said to be in the sample lower half, but for a more accurate localization the thickness must be divi­ded into a higher number of portions. Fig. 3, where sample thick­ness is divided into 8 portions of == I mm, shows that the dela­mination is contained in a material zone between 4 and 6 mm deep.

Similar considerations may be made for samples RI, R8, and R9; their images are not shown to avoid lading of results presentation. A discussion on the quality of the these specimens, based on results obtained with traditional UT methods, may be found in [4 - 5].

The advantages of the complete waveform acquisition technique are not limited to what has been described above. A further capability is represented by the statistical analysis of the material global properties, without resorting to the subjective procedure of image analysis. To survey the statistical properties of the complete waveform database, three-dimensional histograms may be used. A 3-D histogram is formed by color coding the number of occurrences of waveform values in the UT database proportionally to the deviation from the reference waveform [9]. In Fig. 4, the 3-D histograms of the examined samples are reported. A variation of grey tones may be observed in the histograms: the lighter the tone, the higher the deviation from the reference (mode) waveform.

Histograms examination allows for an immediate identification of the global quality of the material by evaluating the width of the deviation and the amplitude of the back echo. Localized defects may be also identified and their position within the thickness accurately defined. Sample Rl histogram (Fig. 4a) is characterized by a low deviation and a high back echo amplitude. This indicates a rather

797

homogeneous structure with negligible porosity, in good agreement with the judgement in Tab. 1. Samples R7 and RS (Figs. 4b and 4c) display a higher deviation and a lower back echo; this suggests a material quality with some diffuse porosity, partially scattering and attenuating the UT energy (Tab. I). Moreover, sample R7 histogram shows a localized high scatter of data indicating a defective zone. It has been previously shown (Fig. 3) that this defect is an extended delamination; its depth was determined by image analysis as comprised between 4 and 6 mm. The localized scatter in the histogram is situated between 4.4 and 5.S mm of depth and its peak value is observed at 5.2 mm; macroscopic examinations indicated a depth between 4.9 and 5.4, confirming the accuracy of histogram measurements. Finally, sample R9 (Fig. 4d) presents a very low deviation and absence of back echo. The material structure seems very uniform but its quality rather low, as UT energy is highly attenuated within the material. This is in good agreement with the high diffuse porosity content of this sample, as reported in Tab. I.

III - CONCLUSIONS

The advantages of an UT NDE technique based on complete waveform digital acquisition and analysis comprise the possibility of obtaining detailed information on the entire material volume, without performing multiple time-consuming C-scans. The localiza­tion of defects within the material thickness is greatly facilitated: in particular, delaminations depth may be accurately measured, provi­ding key information for defect criticality evaluation. Images may be formed from the complete waveform database through standard UT digital procedures and become active interfaces between user and database, allowing for the retrieval of waveforms corresponding to any point of the material scan. Finally, the complete waveform database may be statistically analyzed to produce information on the material global properties.

The disadvantages of the technique are mainly connected with the necessity of resorting to an advanced digital system with very high memory, data processing, graphic, and storage capabilities.

ACKNOWLEOOEMENfS

This work was carried out in part at the Center for Composite Materials (CCM), University of Delaware, Newark, DE 19716, USA. The Italian National Research Council (C.N.R.) is gratefully acknowledged for its support of the present research. Thanks are

798

due to Simpres s.r.1., Lissone (MI), Italy, for providing the experimental materials.

REFERENCES

[1] B. Harris and M. G. Phillips, "Non-Destructive Evaluation of the Quality and Integrity of Reinforced Plastics", Developments in GRP Te c h no log Y - Vol. 1, ed. B. Harris, Applied Science Publishers, London, 1983.

[2] Composite Design Encyclopedia, "Test Methods" - Vol. 6, University of Delaware, Newark, DE 19716, USA, 1984.

[3] R. Teti, "Ultrasonic Inspection of Composite Materials", Proc. J ournees sur Ie Controle Qualite des Materiaux Composites, Bordeaux, 12-13 May 1987.

[4] R. Teti and G. Caprino, "Ultrasonic Testing of Thick GFRP Laminates", Proc. 5th Int. Conf. on Compo Mat., Milan, Italy, 10-12 May 1988.

[5] R. Teti and G. Capri no, "La grafica avanzata nei controlli US computerizzati dei laminati PRFV", Poliplasti, n. 372/3, Nov. -Dec. 1988.

[6] S. A. Weber, "Ultrasonic Non-Destructive Evaluation Based upon Complete Waveform Databases", Master Thesis in Electrical Engineering, University of Delaware, Newark, Del., USA, 1987.

[7] R. A. Blake, "Ultrasonic Image Histogram Evaluation and Enhancement", Proc. 6th Annual Micro-Deleon Conf., IEEE, 1983.

[8] S. A. Weber, R. A. Blake, and R. G. Irwin, "Ultrasonic Image Reconstruction Utilizing a Full Volume Digitized Waveform Database", Review of Progress in Quantitative NDE, Vol. 7B, eds. D. O. Thompson and D. E. Chimenti, Plenum Press, New York, 1987.

[9] S. A. Weber, R. A. Blake, R. Teti, and C. B. Boncelet, "Characterizing Porosity of Composite Materials through Digital Ultrasonic Waveform Processing", Review of Progress in Ouantitative NDE, Vol. 8A, eds. D. O. Thompson and D. E. Chimenti, Plenum Press, New York, 1988.

Tab. 1 - UT samples identification and quality judgement.

Laminate Id Quality Rl Acceptable, negligible porosity R7 Delamination, low diffuse porosity R8 Localized and diffuse porosity R9 High diffuse porosity

799

Fig. 1 - Ultrasonic images of sample R7.

(a)

(b)

(c)

(d)

Fig. 2 - UT waveforms corresponding to the arrow locations in Fig 1.

800

Fig. 3 - UT images obtained by subdividing the thickness of sample R7 into S equal portions.

(Rl)

(R7)

(RS)

(R9)

Fig. 4 - Three-dimensional histograms of the examined samples.

CHOC IMPACT

Chairman: Pr P. LAMiCa Societe Europeenne de Propulsion

INERTIAL EFFECTS IN TWIN SKINNED GRP LAMINATES SUBJECTED TO IMPACT

LOADING IN A THREE POINT BEND CONFIGURATION

ABSTRACf

R.A.W. MINES, C.M. WORRALL, G. GIBSON

University of Liverpool Impact Research Centre - Faculty of Engineering

PO Box 147 - L69 3BX LIVERPOOL - England

Results are given for impact tests for a mass of 4.3 kg dropped from a height of up to 3m onto a GRP sandwich beam in a three point bend configuration. Force versus deflection traces are derived from accelerometer data and the effect of span and impact velocity discussed. A two degree of freedom inertial model is developed for the 180 mm span test to give the local deformation of specimen in the vicinity of the impactor and hence the effect of local stresses on failure. High speed photographic data is used to supplement experimental results.

1. INTRODUCflON

Twin skin construction with composites offers the possibility of achieving a considerable increase in the rigidity of beams and panels with very little increase in weight. Although the principles of static design of sandwich panels are well known 111 the same is not yet true where design for impact is concerned. Composite sandwich panels and beams can be subject to impact by objects with a wide range of speeds, depending on application. Often impacts in the relatively low speed range of a few metres per second can be critical. This corresponds to low velocity collision and 'dropped object' damage.

The laminate selected for study consists of a non-woven polyester core, know as Coremat, with woven roving and chopped strand mat reinforced skins. Polyester resin was used as the matrix for the skin and core. Such construction is widely used in marine and land transport applications. Beam specimens were tested under simply supported three point bend conditions with central impact. Some preliminary results have already been quoted in the literature 121, in which a 4.7 kg mass was dropped from a height of up to 1 m and results derived using a Laser-Doppler velocirneter.

804

2. EXPERIMENTAL

The sandwich laminates from which the bend specimens were cut were prepared by hand laminating and contact moulding. The skins each consisted of one ply of 360 GSM plain weave woven roving and one ply of 450 GSM chopped strand mat, the latter being on the surface of the specimen. This construction is different to that of the tests quoted in 121. A typical skin thickness was O.Smm. The cores consisted of two layers of 5mm (nominal) Coremat 131, a propriety core material of non-woven organic fibre containing hollow microspheres. The effective core thickness after moulding was S.7mm, the total section thickness being 1O.3mm. The polyester resin was Scott Bader Crystic 272.

The impact tests were achieved using a 3m high drop hammer. Details of the impact area and specimen are shown in figure 1. Two beam spans were selected for study, namely ISOmm and 400mm. It was ascertained by making trial tests at different span to depth ratios in the static configuration that, over the elastic portion of the load deflection curve, shear deflection accounted for 12.S% of the total deflection for the ISOmm span specimen and 2.6% for the 400mm span specimen.

The deceleration of the falling mass through a test was measured by an accelerometer mounted in the impactor nose and with a frequency response of 50kHz. The signal was amplified using a Tektronix IMHz differential amplifier and stored on a 20MHz DLlOSO Datalab transient recorder. The signal was then filtered using a digital analogue of an R-C filter, with a cut off frequency of 500Hz. The signal was then processed to derive impact mass applied force and displacement

To supplement accelerometer data, some high speed cine films were taken of selected impact tests. A Hadland Hispeed camera was run at a nominal 7000 fps giving an interframe time of 143J,1s. The camera was focussed on the side of the specimens, in the vicinity of the impactor. Typically the field of view was l00mm x l00mm.

3. ISOmm SPAN TESTS

These tests were similar to those quoted in 121, but with an improved design of impact mass. Force v deflection results were similar to those quoted in 121, which were derived using a Laser Doppler Velocimeter. A typical result is shown in figure 2. The deflection at beam failure as predicted from high speed photography agrees closely with the maximum load accelerometer results. Of the three tests conducted at a drop height of l.5m, two failed as a result of core shearing and one by upper skin buckling (see figure 2). The force v deflection traces were similar up to beam failure and hence it was concluded that failure could result from either bending or shear dominated behaviour. This type of behaviour has been found for static loading of sandwich beams 141 in which seven possible modes of failure can occur. In order to ensure that bending effects dominated, it was decided to increase the beam span to 400mm.

Of interest in impact tests is the separation of structural inertia effects from material strain rate effects. A simple model that has been proposed for quantifying the inertial behaviour of three point bend tests 15,61 consists of a two degree of freedom mass spring model in which the impact mass, effective beam mass, contact stiffness and beam stiffness are modelled. The model and values of physical parameters are shown in figure 3. The contact stiffness k was measured by placing a specimen on a flat rigid surface and slowly driving the impactor into it. The beam effective mass assumes small deflection linear elastic response and the specimen stiffness assumes that

805

the fundamental mode of vibration dominates. A result from this model for a drop height of l.5m is shown in figure 2. Some viscous damping was included in the beam stiffness to ensure a decay of inertial oscillations. The two coupled second order linear differential equations of motion were solved using a 4th Order Runge Kutta numerical algorithm. The numerical model shows high frequency (2.2 kHz) oscillations, due to contact stiffness effects, superimposed on the fundamental mode stiffness of the beam. This high frequency was not found in the accelerometer results, due to the 500Hz filtering, but were suggested by the high speed photography results (see figure 4). Figure 4 predicts a contact stiffness of 6.9 x lot> Nm-l and shows that the impactor indentation at failure is approximately half a millimeter.

Bending dominated failure (figure 2, Mode 1) could be influenced by global bending of upper skin or by local indentation stresses. It is proposed that the latter would become important at high impact velocities. A simple strength of materials model was developed in which global bending was modelled using standard beam theory 111 and local effects were modelled as a cylindrical penetrator defonmng the upper skin on an elastic foundation, i.e. the Coremat core. Such an analysis predicts, for the deflection at failure met here, that the ratio of global to local compressive stresses in the upper surface of the upper skin were in a ratio of 1 :2. This ratio depends on impactor penetration, radius of impactor and on material properties.

4. 400mm SPAN TESTS.

The span of the three point bend configuration specimen was increased to 400mm to ensure that bending effects dominated. Fifteen tests were conducted, three at each drop height from 1'2m to 21/2m in half metre steps. Figure 5 shows accelerometer results for a drop height of 1m and 2m. In figure 5 the structural response is dominated by an oscillation of frequency 400Hz which is not associated with contact stiffness or fundamental beam vibration mode effects. Such oscillations could be due to higher vibration modes or flexural waves travelling along the beam. For the drop height of 1m the high speed cinE film gives upper skin failure at maximum load and also gives a loss of contact. Loss of contact effects have been observed in other three point bend tests 171. In the case of drop height of 2m failure is shown to occur at the peak after the maximum load suggesting an effect of incremental damage. Also two major losses of contact are shown in the films. The failure of the upper skin was less localised than in the case of the l80mm span, suggesting that stresses in the vicinity of the impactor were dominated by global effects.

Inertial analysis of the 400mm span test is more complex than in the l80mm span case. Loss of contact makes a simple two degree of freedom inertial model inapplicable. The specimen fails at a time which is typically much less than the fundamental period of vibration of the specimen hence an effective mass and global specimen stiffness become inappropriate. Large deflections, typically four times the specimen thickness, make a small deflection linear elastic theory inapplicable. Hence a more sophisticated inertial model is required in this case. The deflections at beam failure from accelerometer and high speed photographic data were: static - 54mm, 1m drop height - 39mm, 11/2m - 37mm, 2m - 42mm, 21/2m - 48mm.

5. CONCLUSIONS

The failure of GRP sandwich beams, of a specific thickness, under impact

806

loading is dependent on impactor mass, impactor geometry, impactor velocity, beam span and on boundary conditions, i.e. clamped or simply supported. It has been shown that the mode of beam failure is dependent of the relative values of bending and shear. Failure is also influenced by local and global stress levels. It has been shown that the inertial behaviour of the test is also dependent on the above effects. It is proposed that simple inertial models should be developed, which portray the gross behaviour of the test, and which can be adapted to the different impact conditions. In this way such models could be used as predictive models for a given set of impact conditions.

ACKNOWlEDGEMENTS

We are grateful to Mr. K. Godber, GRP Material Supplies Ltd and Pilkingtons PLC for the supply of the reinforcements and Coremat cores which were used in this work. CMW is supported by a SERC studentship. Dr. R.S. Birch of the Impact Research Centre assisted in the dynamic tests and wrote the computer analysis software.

REFERENCES

1. H.G. Allen, "Analysis and Design of Structural Sandwich Panels" (1969) Pergamon Press.

2. A.G. Gibson, R.A.W. Mines, C.M. Worrall in "Fibre Reinforced Composites 1988" (Ed. A.G. Gibson) (1988),20.1-20.9.

3. "Shaping the Future", Coremat product catalogue, GRP Materials Supplies Ltd., Alchorne Place, Burfields, Portsmouth P03 5QU.

4. T.C. Triantafillou, L.J. Gibson, Materials Science and Engineering, Vol. 95 (1987), 37-53.

5. W. Suaris, S.P. Shah, Cement, Concrete and Aggregates, Vol. 3 No.2 (1981), 77-83.

6. J.G. Williams, G.c. Adams, International Journal of Fracture, Vol. 33 (1987), 209-222.

7. W. Bohme, J. Kalthoff, International Journal of Fracture, Vol. 20 (1982), 139-193.

mm.

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m, :4Bkg

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[, : 0

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Figure 3 - Two degree of freedom model with parameters.

808

Local Penetration -

160 ,,1ft Span "5", Drop Height

Stresses In Upper Slun .-

Global Local Total

Figure 4 - Indentation values from high speed photography (HSP) results, with a simple local stress model.

Span = 400mm . Filter Frequency: 500Hz. Drop Height = , m

1500 \Repeat"Tes:,

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Figure 5 - Force v deflection from accelerometer results for drop heights of 1m and 2m for 400mm span tests (LOC - Loss of Contact)

ABSTRACT

THE EFFECT OF CRYSTALLINITY ON THE IMPACT PROPERTIES OF ADVANCED

THERMOPLASTIC COMPOSITES

D. LEICY. P.J. HOGG

Queen Mary College - Department of Materials Mile End Road - E1 4NS LONDON - England

A range of thennoplastic matrix carbon fibre composites with varying degrees of crystallinity has been subjected to falling weight impact tests . Tests perfonned at excess energies. sufficient to cause total penetration. indicate that the thennoplastic composites are superior to toughened epoxies in tenns of energy absorbed. After low energy impacts. the thermoplastic composites absorb less energy then epoxies with this difference being attributed to a reduction in the amount of microcracking. PEEK based composites are shown to be superior to PPS based composites and amorphous systems superior to crystalline equivalents. The generation of a plastic dent on the surface of a thennoplastic composite at an early stage of the impact event is thought to modify the deformation behaviour of thennoplastics relative to thennosets .

1. INTRODUCTION

Many high performance carbon fibre composites are now available with thennoplastic matrices as an alternative to the more traditional thermosetting matrices typically based on epoxy resins. The new thennoplastics are claimed to offer certain advantages notably with respect to manufacturing costs compared to thennosets. with improved toughness and environmental resistance providing additional benefits [1]. The need to provide environmental resistance. notably in aggressive aerospace fluids Get fuel. paint strippers. de-icing fluids etc ) has favoured the introduction of semi-crystalline thennoplastics such as PEEK and PPS over amorphous materials such as PES [2]. The use of semi-crystalline matrices complicates the manufacturing process by introducing an additional morphological variable. such that the structure will to some extent be dependant on thennal history . Although much effort has been expended by the raw materials suppliers in developing matrices that are tolerant of variations in cooling rate • the degree of crystallinity will be controlled by the rate of cooling and in extreme situations the development of crystallinity may be suppressed [3,4]. A key property for advanced composites for use in aerospace applications is the impact resistance of the material [5]. It is generally accepted that the new generation of thennoPlastic matrix carbon­fibre composites are superior to conventional epoxy resin laminates although the basis for the improved performance is not established other than through the implicit assumption that the

810

matrix is tougher. It is also accepted that the presence of crystalline regions in a thermoplastic matrix will be detrimental to mechanical properties such as toughness by acting in an analogous fashion to cross-links. The work reported in this paper was concerned with establishing the mechanisms by which thermoplastic composites exhibited superior impact performance to thermosets, and simultaneously investigating the effects of extreme variations in processing and therefore crystallinity, on this impact behaviour. The major emphasis was placed on the consequences of relatively low energy, non-penetrating impact blows on composite laminates rather than excess energy, fully penetrating impacts. The former situation mirrors the practical issues of concern to the aerospace industry and is the initial step in the critical .. damage tolerance "test which has been adopted in various guises as a requirement for certification of a composite material.

2. TEST METIIODS AND MATERIALS

The carbon fibre reinforced composites used in this programme all utilised Hercules AS4 fibres in combination with either PEEK (APC-2 from ICI) or PPS (Ryton from Phillips 66). All materials tested were produced as quasi-isotropic laminates with a stacking sequence [ 0,+45,90,-45 ]sym .Nominal thicknesses of Imm, 2mm, 3mm and 4mm were prepared by using 8, 16, 24 and 32 plies of prepreg. Laminates were either allowed to cool according to manufacturers recommendations after processing at the melting temperature or rapidly quenched. These conditions produced two extremes in the measured degree of crystallinity in the matrices. The PEEK matrix was found to have a measured ( by DSC) degree of crystallinity of 35% when cooled normally, which fell to 8% after quenching. The APC-2 prepreg had a fibre volume fraction of 61 %. The PPS matrix possessed a degree of crystallinity of 50% under standard processing conditions and 5% after quench cooling. The fibre volume fraction in this case was 53% Throughout this paper, the material cooled under standard conditions (both PEEK and PPS ) is referred to as crystalline , and the material that has been quenched cooled is referred to as amorphous. This is despite the fact that it is recognised that both systems are in reality semi­crystalline. A limited set of data has been obtained from other workers on unidirectional thermoplastic composites processed in an identical fashion and on toughened epoxy laminates incorporating identical fibres. The epoxy resin was not identified in the latter case. Mechanical testing involved a series of short beam shear tests performed according to ASTM in order to characterise and compare materials, and a series of high and low energy instrumented falling weight impact tests. The instrumented impact tests were performed on two machines equipped with identical software and instrumentation. One machine was a proprietory CEAST drop-weight tower used in conjunction with an AFS mk3 Fracturescope and this was used for high energy full penetration tests. The second machine was built in-house and featured a lower mass carriage and striker system suitable for lower energy, non-penetration tests together with an anti-rebound device to prevent multiple impacts. On both machines the striker consisted of a steel tup with a 20mm hemispherical tip and the specimen was simply supported on a steel ring, diameter 40 mm. All impact specimens were flat plates 60mm x 60 mm. The raw test data comprised a force-time record which was subsequently processed to provide information on force, deflection, velocity, and energy absorbed at any interval during the test. Impact tests were performed within the velocity range 1-4 mis, depending on, the incident energy required. No rate effect was observed during testing within this range and impact results quoted refer to incident energy and not impact velocity.

3. RESULTS

The results of excess energy impacts are shown in figure 1 for specimens of thickness 2mm. It is clear that the PEEK matrix systems are superior to the PPS matrix composites with a maximum energy absorption in the range of 42-52 J compared to 25-35 J. Some of this difference may be ascribed to the lower fibre volume fraction in the PPS composites but this would not fully account for such a large disparity in the results. A representative value for a toughened epoxy system would be of the order of 301. There does not appear to be a consistent trend in these results as a result of crystallinity.

811

The behaviour of the various systems at low energies is somewhat different. Figure 2 shows the energy absorbed by the various laminates after low energy impacts. In this case it is the PPS composites that exhibit the greater energy absorption of the thermoplastic composites and data obtained for the epoxy system shows it to absorb relatively similar energy to the PPS laminates. A general trend is observed if the absorbed energies and peak forces generated during the low energy impacts are considered for all materials over a range of thicknesses. The absorbed energy increases with increasing impact incident energy, with the PPS composites absorbing more energy than the PEEK composites and with the crystalline laminates absorbing more energy than the amorphous laminates, figure 3. Similarly, the peak forces increase with increasing impact incident energy, with the PEEK laminates exhibiting a higher force than the PPS laminates and with the crystalline composites exhibiting a lower peak force than the amorphous composites, figure 3. Figure 4 illustrates the trends in peak force versus incident impact energy fof 24 ply (3mm thick) specimens and includes data obtained by co-worters for a comparable epoxy system. Comparing figure 2 and 4 will show that the epoxy system exhibits similar behaviour to the PPS matrix laminates in terms of energy absorption (high) but simultaneously is similar to the PEEK matrix laminates in terms of peak force (high). Clearly there are fundamental differences in the deformation of the thermosetting and thermoplastic matrix composites. An important difference between thermoplastic andthennosetting laminates that was observed here and has been reported elsewhere is the formation of a small indentation on the surface of thermoplastic composites after very low energy impacts. The deformation behaviour of thermoplastic laminates is influenced by the generation of this dent on the impacted surface which progressively leads to the formation of a dome on the tensile surface of the specimen at increasing incident energies. Catastrophic fracture in the thermoplastic specimen is believed to proceed by' a relatively stable tearing across the dome prior to large scale fracture. In thermosetting systems by contrast, cracking initiates much earlier on the tensile surface with no attendant dent on the impacted face. The fracture process proceeds by a series of sudden bursts of splitting and delaminations and is generally more unstable in character [6]. Recent work has shown that the onset of damage in carbon fibre composites may be linked to the position of a sub-peak on the force deflection curve and that this sub-peak occurs somewhat later (or requires an impact blow of greater energy) for the thermoplastic PEEK matrix composite than a comparable epoxy [7]. The sub-peak is definitely linked to the onset of cracking and does not indicate yielding of the matrix to form a dent .This is illustrated by the series of schematic micrographs shown in figure 5 alongside the relevant force-time curves.

4. EFFECTS OF CRYSTALLINITY ON DAMAGE

The nature of the cracking pattern in a laminate is to some extent governed by the thickness and therefore stiffness of the, specimen. Thin specimens tend to deform extensively under the impactor and cracking initiates as transverse cracks near or at the tensile surface [8]. These cracks subsequently connect via a network of delaminations and shear cracks. If the laminate is thick and stiff, then deflection under the impactor is reduced but higher contact stress are generated and fracture initiates as compressive cracks under the impactor itself [9]. Once again these cracks propagate and connect via delaminations and shear cracks. The exact thickness required to define a transition point from one mode of failure to another will be dependant on the strengths of the material under the specific loading conditions. Crystalline PEEK within the range of thicknesses tested always failed in flexure with cracks originating at or near the tensile face. The effect of producing an amorphous equivalent material was to delay the onset of cracking. By way of an example, for a 24 ply specimen, small transverse cracks were first detected in the crystalline material after low energy blows of 2.0 1 incident energy. No damage was observed in the the amorphous laminate after such a blow and a blow of 4.21 was required to produce a similar level of cracking. Figure 6 compares the crystalline and amorphous laminates after blows of 4.21. The behaviour of the PPS laminates was different. Amorphous PPS laminates initially show cracking near the impacted face, and this after comparatively low energy blows - eg 0.981 for a 24 ply specimen. A 4.21 blow inflicts considerable damage on a PPS laminate, Figure 6, that is more extensive than that observed in the crystalline PEEK specimens. The crystalline PPS specimens however show even more severe cracking. In this case however, the cracking occurs initially along the centre of the laminates and consists of interlaminar cracks presumably resulting from

812

the induced interlaminar shear stresses. This behaviour is typified by the schematic in figure 6 showing the damage after a 4.2 I blow on a 24 ply specimen. The results of the interlaminar shear tests performed on both quasi -isotropic specimens and external results on unidirectional equivalent material reveals a trend in properties that match those observed in the cracking behaviour, table 1. The transverse tensile strength of amorphous PEEK laminates is higher than that of the crystalline PEEK which is itself higher than that of the amorphous PPS laminates and the interlaminar shear strength of crystalline PPS (quasi -isotropic) is exceptionally low.

5. DISCUSSION

The results indicate that the performance of a particular composite is dependant on the mechanism of failure and in particular the nature of the initiation of damage. Epoxy resin laminates appear to be good at absorbing energy after low energy impacts because they are readily susceptible to cracking at low energies. The nature of the cracking is such however that the stresses in the laminate continue to rise rapidly and this has the consequence of producing an unstable fracture process characterised by large, sudden bursts of splitting. The total energy absorbing capabilities of the epoxy laminates are limited by this process. As a general class of materials, thermoplastics would appear to be more efficient overall in absorbing energy. With these materials the absorption of energy again appears to be clearly linked to the formation of cracks rather than by any yielding processes in the matrix. The improvements relative to thermosets in energy absorption after total penetration are probably linked to the development of the initial indentation which would have the effect of reducing the local stresses under the impactor and allowing a greater density of cracking to develop in a stable fashion. These differences in deformation behaviour are a possible explanation for the inconsistent trends shown by the epoxy data in figures 2 and 4 compared with the· thermoplastics. Within the thermoplastic composites as a class, the general rules were that the greater the extent of cracking mat occurred, then the lower the peak force generated and the greater the energy absorbed. The susceptibility of the laminates to damage under impact conditions would appear to be indicated by the simple mechanical property measurements obtained at conventional testing speeds. Similarly the effects of processing on impact properties are reflected by the effects of processing on these simple properties i.e. transverse tensile strength and interlaminar shear strength . The low energy absorption of the PEEK laminates, and particularly amorphous PEEK laminates after low incident energy impacts is the result of a resistance to cracking within these materials. This suggests that such materials would also exhibit a superior damage tolerance as any post impact destructive test would be critically dependant on the levels of damage within the specimen.

6. CONCLUSIONS

The mechanisms of failure under out-of-plane impact conditions are different for thermosetting and thermoplastic composites. It is likely that the differences are responsible for the increased energy absorption in thermoplastic composites, especially those base on PEEK, during a total penetration impact. The degree of crystallinity will affect both the impact behaviour of a laminate and the basic laminate properties such as interlaminar shear strength and transverse tensile strength. Energy absorption after low energy impacts is a reflection of the amount of cracking within the material.

ACKNOWLEDGEMENTS The cooperation of Dr A Bunsell and Ms Caroline Truffier at the Ecole de Mines, Paris in processing and supplying material and test results on unidirectional material is warmly appreciated. Discussions with our colleagues Dr S Turner and Mr C Dunn at Queen Mary College have also proved most helpful .

REFERENCES 1. D.C. Leach and D.R. Moore, Compo Sci. Tech. 23. (1985) 131-161. 2. S. Witzler, Advanced Composites, March/April (1988) 55-60 3. C.M. Tung and P J. Dynes, I. Applied Polymer Science, 33 (1987) 505-520.

4. P. Cebe, S.Y. Chung and S. Hong, J. Applied Polymer Science, 33 (1987) 487-503. 5. S.M. Bishop, Composite Stuctures, 3 (1985) 295-318. 6. P J. Hogg, C. Dunn and A. Ahmadnia, Unpublished Information, (Queen Mary College,

London, UK.) 7. P.E. Reed and S. Turner, Composites, Vol 19, No.3 May (1988) 193-203. 8. W J. Cantwell and J. Morton, Composite Stuctures, 3 (1985) 241-257. 9. W.J. Cantwell, PhD. Thesis, University of London, (1985).

TABLE 1

Carbor Fibre Orientation Unidirectional Quasi-IsotroPic Composite Materials CFPEEK CFPPS CFPEEK Processing Conditions AM./CR. AM./CR. AM./CR. Flexural Stren~ (MPa) 1200/1800 825/1000 815/892 Interlaminar Shear Strengtlt{MPa) 65.0/95.0 46.0/38.0 56.0/77.0 Transverse Tensile Stren2th (MPa' 85.0/60.0 30.0/20.0 ••••

L-~~~~~~i~~e~'~e~'~6~'~~'~~~: O

.... o~~~~9 0

f i \ I

(f) 038110S8V ADlI3:N3:

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34.0/16.0 ••••

813

814

8 PLY LAMINATES

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5 3

.98 2 4.2 IMPACT ENERGY (J)

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FIGURE 3: THE RELATIONSHIP BETWEEN PEAK FORCE & IMPACT ENERGY

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815

Figure 5: Onset of damage in 24 ply crystalline PEEK demonstrated by the first peak on Force·time curves

Figure 6: schematic damage development or 24 ply crystalline and amorphous thermoplastic composites subjected to low velocity impact (apparent damage immediately under impactor due to indentation, not cracking)

--==~=-=- --=------=-= -------

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Communication parvenue hors d61ai Late paper

817

ENDOMMAGEMENT DE STRUCTURES TUBULAIRES COMPOSITES

SOUS SOLLICITATIONS DYNAMIQUES

P. HAMELIN, C. BURTIN

CERMAC -INSA LYON Bat 304 - 69621 VILLEURBANNE Cedex - France

'!he aim of this paper is to stmy the shock behaviour of ClCIIpOSite tubes mDer oatpreSSive Ioad:in;J cxnli.tioos. We prove that the type am nature of the fibers, the mec:banical prcperties of the matrix, the gecmetry of the structure am the fiber arran;JE!JD9l1t affect significantly the energy absoIptim capabilities of the cx:up:lSite tube.

Ie prdJIeme d' absorptim d' energie et de resistance au choc fait cq:para1tre de J'lCIIi>reux ~ lies a Ia nature het:erogene des materiaux ClCIIpOSites inpliquant des ~ CXIIplexes lies a Ia nature het:erogene des materiaux ClCIIpOSites inpliquant des ~ de rupture au niveau de Ia matrioe, de Ia fibre, de I 'interface, de cIelaminage, de CXII'binaisals de divers processus. Act:uellement la repcnse des materiaux ClCIIpOSites a une sollicitation dynamique est eomee de maniere semi-enpirique. Cette ~ est rendue difficile par Ie l'XJIbre eleve de cx:IIbi.nai.sals des materiaux oonstituants, des sequeooes d'ESlpilement et par Ies effets de oarportEment struct:uraux 00 Ies :interacti.oos entre oarportEment dEpen:)ant des materiaux et repcnse en defOImation des structures. Noos pc:uva1S citer Ies travaux de 'lbomta1 /1/, Faye /2/, Kimervater /3/, crcnkhi.te /4/, Provensal /5/, Il1lI /6/ qui abordent de fClQOll essentiellement ~Iogique Ie prdJIeme.

820

I - IKlCESSUS D' ABSORPl'I~ D'mERiIE

1.1. Absorption d' energie

I.es tubes cirallaires a matrice thenoodurcissab1e et a renforts fibreux presentent pr:irx:ipale.ment deux types de OCIIPOrtement en catpCession axiale (quasi-statique) :

- tans Ie premier JOOde (JOOde I) (Fig. 1), 1es tubes a ba.rt:s plats peuvent al1er a la ruine de deux manieres : a faible efD;z ils se cassent par instabilite en flanilage et a fort efD;z par depassement des ccntraintes admissibles du materiau saJS sollicitation de CCIIpression. Ceci cxn:1uit a de faibles valeurs de Wsc avec preserre d '\.Dl pic Fer tres inportant. II en r8sul.te une ruine de 1a structure dans son enseni:l1e. la plus grame partie de l' energie absomee se situe dans la P'lase de OCIIPOrtement elastique.

- rans Ie deuxieme JOOde (lOOde II) (Fig. 1), 1 'ecrasement progresse a partir du boot du tube en contact avec Ie plateau JOObile de 1a presse. rans 1a zone d'ecrasement, Ie tube passe d'\.Dl etat non erDmnage a \.Dl etat CCIIpOrtant une lIIlltitude de ''microfissures''. Cela ne signifie pas qu I il ait une faible resi.stan::e a la prcpagation des microfissures, car il y a une lIIlltitude de microfractures susoeptibles d I absorber de 11 energie elastique mise en cause par 1a diffusion des "cracks".

L' energie d I absorption deperrl dor'x: direct:ement de 1a force myenne d'ecrasement Fm et est irrlepen:iante de la IC>nJUeUr du tube.

1. 2. Mecanisme d I initiation

I.es JOOdes I et II de rupture peuvent etre consideres camne corx::urre.nti.els parce que. 11 initiation de 11\.Dl des deux exc1ue I 'autre. On abtient 11\.Dl des deux suivant 1a dlarge Fer necessaire pour initier une premiere rupture. Poor deux tubes identiques, Ie deuxieme m::x:ie necessite une charge mins :inp:>rtante que pour Ie premier d I 00. 1'idee de favoriser celui -ci par 1a reaJ.isation d 1\.Dl defaut initial : noos usinons Ie boot du tube 00. sera awliquee 1a dlarge en reaJ.isant \.Dl d1anfrein a 45· sur son pourtoor. la relation charge-c:iep1acement est DDdifiee suivant Ie pr:irx:ipe decrit par 1a figure 2.

II - PROGRAMME EXPERIMENl'AL

2 .1. Materiaux et proc::Edes

la matrice est constituee soit de resine Viny1ester Derakane c:u de resine Epoxy B8. I.es renforts enp10yes se presentent saJS fonne de stratifil c:u de tissu ~ de fibres de Verre E, de fibres de cartxme, de fibres de Kevlar, c:u d'\.Dl ~ hybride avec des fibres de cartxme et une gaine en Kev1ar. I.es choix des proc::Edes de fabrication tient ccmpte de la possibilit.e d 'utiliser des techniques semi -irrlustriel1es adaptees a 1a reaJ.isation de corps creux : nnllage par enroolement filamentaire, nnllage par errluction de resine sur des tresses confol.'Illl3es sur \.Dl mamrin.

821

2.2. CcJipositioo des tubes utilises

Noos etu:iians plusieurs familIes de materiaux c::x:mpcsites. res parametres intervenant dans la fOl:1ll.llatioo des materiaux et la fabricatioo soot :resumes dans Ie tableau ci -dessaJS :

serie Resina Fibre TF &!pil. Olanf. +8 traJs

1 Vinyl 470-36 Verre E EF (±45) 3 em 2 Vinyl 470-36 Verre E EF (±45) 3 cui cui 3 ~ IER332 Verre E 'lR (±52) 4 cui 4 ~ XBS084 Verre E 'lR (±52) 4 em 5 ~ IER331 Verre E 'lR (±52) 6 cui 6 ~ IER331 Verre E 'lR (±52) 4 cui 7 ~B8 cartlale TJOO EF (±45) 8 cui 8 ~B8 carl:la'lefKevlar EF (±45) 6 cui 9 ~B8 Kevlar 49 EF (±45) 4 cui cui

Caract:eristiques des tubes enployes.

avec: TF : tedmique de fabricatioo, EF : enroolement filamentaire, 'lR : tresses. Toos les stratifies soot antisymetriques.

2.3. Dispositif d'essais

Noos utilisons W1 canon a air OCIlprime prcpllsant W1 ballet de 100 kg a des vitesses entre 0 et 20 nVs. la mise au point du dispositif ainsi que la definitioo de la procEdure exper:iJnentale a ete realise par Hamelin, Rlria et autin /8/.

3 - RESUIlI'ATS - SYNlHESE

3.1. Modes de rupture

Des eprcuvettes tub.llaires en materiaux oatpOSites avec W1 eJ.anoement CCIIpris entre 7 et 10 oot em inpa.ct.ees a des vitesses de sollicitatioo CCIIprises entre 16 lis) et 38 (l/s) et des energies cinetiques CCIIprises entre 3,4 kJ et 19,5 kJ. I.e lOOde de rupture est de nature instable et non reproductif poor des tubes dant les extremites soot planes. Un mecanisme geaoetrique (W1 dlanfrein) jouant Ie role de c:orx:entrateur de contraintes et situe a 1 'extremite de la structure est neoessaire poor passer d'W1 lOOde d'9crasement instable (lOOde I) a W1 m:x:1e stable et reproductif d'9crasement (m:x:1e II). I.e m:x:le II d'9crasement est conservatif dans le danaine de vitesses etu:iie (entre 5 et 10 nVs). I.e lOOde d'9crasement stable deperrl a la fois de la nature de la fibre utilisee et du precEde de fabricatioo retenu. res types de rupture stables oorrespadent a :

- des coodles qui s'evasent ecrase.nent type "fleur" (canposites a renfort fibres de verre E/enroolement filamentaire).

- des coodles qui flecnissent : 9crasement type "dlaussette" (canposites a renfort fibres de verre E saJS fonne de tresses) .

- des microfragmentatians (c::x:mpcsites a renfort fibres de

822

carlxlne CAl carl:ale-Kevlar/pI'OOEde d'enrcW.ement filanart:ai.re).

Ia cc:atparaison entre Ie c:x:atpOrtaDent statique et au dloc est etablie pcm' les differentes series (Fig. 3). I.es series 1 et 2 (renfort fibres de verre Ejmatrioe Vinylester 47o-36/enrcW.ement filanart:ai.re) presentent 1m JOOde de zupture stable (JOOde II) en regilne quasi -statique et sous sollicitatioo de dloc. si les tubes ne possedent pas d' initiateur, Ie JOOde de zupture est instable (JOOde I) .

En regilne statique les tubes (series 3 a 9) a base de fibres de vene E sous forme de tresses avec une matrioe :EpJxy CAl de fibres de carlxlne, carl:x:1le,IKevlar et de fibres de Kevlar (pI'OOEde d'enrcW.ement filamentaire) avec une matrioe :EpJxy tendent vers 1m ocuportement. de type instable different du JOOde stable dJt:eru sous sollicitatioo de dloc. L'initiateur en statique ne jooe plus sal role d'aJlDrceur de zupture progressive. I.es cootraintes maxima satt confax:lues avec les contraintes critiques en statique et au dloc pcm' ces series (4-5-6-7-8-9) . Ia fomatioo d tune coormne au del:ut de I' essai est neoessaire pcm' aJlDrcer oet ecrasement progressif pcm' les series 3 a 6 (renfort fibres de verre Ejtresses). Poor la serie 9 Ie c:arportement est identique a oelui ciJserve sous sollicitatioo de c:hoc et peut etre c:arpare a 1m flanmge localise voisin de oelui rerx::ontre au sein de tubes metal.liques. Ce type de flambage est la consE!querx:e du caractere ductile de la fibre de Kevlar.

3.2. Influence de la vitesse de sollicitation sur l'energie specifique

Un effet de dynamicite est dJt:eru lorsque la vitesse de sollicitation croit dep.ris Ie reg.ilne statique jusqu'a la vitesse de 19 DVs se traduisant par une augmentatioo ocntinue des valeurs des contraintes lIDYennes, maxllrums et des energies specifiques. Ia majoration a awliquer deperrl de la serie cmsideree a une vitesse cIonnee. Il existe 1m seuil energetique oon:iuisant a I' erxlcmnagement des tubes. Celui-ci est oarpris entre 1000 et 3000 J. Poor 60 % d'ecrasement, JnlS et:ablissoos 1m classement entre les series.

serie Lv'Iu = 60 % Wsc kJ,Ikg serie Lv'Iu = 60 % Wsc kJ,Ikg

5 41,6 ms 67 4 51,8 ms 57 3 48,0 ms 58 7 39,1 ms 55 8 42,5 ms 57 2 54,6 ms 32 lb 62,7 ms 27

Energie specifique et cluree du c:hoc pcm' 1m rawort Lv'Iu = 60 %.

L'absorptioo d'energie specifique a~ lorsque Ie rawort 4aisseur sur cliamet:re (efD2) croit pcm' des tubes CXJnStitues des memes materiaux et en reg.ilne de dloc.

823

Noos pouvalS retenir que Ie OCIIpOrtement au dloc des mat.eriaux et des st::ructures OCIIp:lSites, analyse en fcn::ticn de parametres tels que 1 'energie al::lsoxbee, les cxntraintes naninales, les tassements suwortes est superieur a Oel.ui de st::ructures en acier au en alum!nim. Toute la difficul.te pour valariser leur awlicaticn vis­a""Vis de oe type d' awlicaticn parte sur 1 'qtimisaticn de la fOl:1ll1l.aticn et du dimensicnnaaent de oes JlBt:eriaux en fcn::ticn des JXIIbreux parametres (nature des .tari.aux, fabricaticn ••• ). Notre etude JID1b:e que, pour les gaDIES de vitesses de chargement retenJes, les diff8rerx:les entre CXIIp)rt:aIIent staticpe et dloc des tubes OCIIp:lSites depenient. des lois de 0CIIpll':tement des mat:eriaux OCIlStitutifs.

FCR FORCE (ItN) 40

( 2)

, ~ 30

(2)

20

10

(1)

Fig.1 - Relaticn c:haIge-deplaoement en OC'IIpresSicn

\ \ ,

TUBE SANS CHANFREIN

TUBE AVEC CHANFREIN

\" --------l~·

CHARGE D'ECRASEMENT MOYENNE

DEPLACEMENT

Fig.2 - Influence du dlanfrein sur 1a relatioo ~1acement

824

NUI.IEROS DES SERIES

IS::S:] I./D.g ---;;

!~ V_ t • .wn/. STATIQUE I >< ,-

rz::2]I./O. 4 ,5 I :> CHOC V_1 2.8m/a

-, -V~I ""'/. 60 , I , , ,-, < I ,

'"' '" SO 0<

~ " '" ~ , w .,

" , J .0 I 0 : I

i: I

V , W I "- 30 II>

W

" '" w 20 z

w

10 /

0 .,

3 4 6 7 8 9

Fig. 3 - Ccl!paraisal entre l' energie specifique en statique et au dloc

1 - 'lhomton P.H., Energy abscnptien in OCIIpOSite structures, Joomal of OCIIpOSite materials (1977), V.13, 241-262

2 - Faye R.L., A crashworthiness test for OCIIpOSite fuselage sttucture, Fibroos OCIIpOSite in structural design, New-York, plennn press, (1980), 214-257

3 - :K:imeJ:vater, Energy absorbirg qualities of fiber reinforced plastic tubes, Prooeedirgs of the american helicqrt:er national specialists meetirg in CCIIp)Site structures, RrlladelIirla Pennsylvania, 23-25 mars (1983)

4 Cronkhite J. D. , Investigatien of the crash, Inpact dlaracteristics of advarn:!d airframe structures, J. of the american heliccpter society, Octd::>re (1982), V. 26, n° 4, 52-163

5 - Provensal, Dissipatien d' energie dans les structures en materiaux OCIIpOSites, Prcx:leedin:j of the 18th intenlational FISIAA <::on:Jress, 3-8 mai (1980), 70-76

6 - Hull D., Effect of speed en progressive CI'UShin;J of epoxy glass cloth-tubes, 3rd

7 - MJria Vila D., cootribltien a 1 'etOOe du cx:q:x>rtement au choc des batons, 'Ibese de doctorat, Institut National des Sciences AWliquees de LYCfi, LYCfi I, (1986), 300 P

8 - Burtin C., Hanelin P., Crash inpact behavior of tubllar CCIIp)Site structures, Int. Conf. en inpact loadirg arxl dynamic behaviour of materials, Breroc>n, RFA, 18-22 mai (1987)

COMMITTEE E.A.C.M.

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Comire d'Expansion Aquitaine

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F.x. de CHARENTENA Y, P .LAMICQ, R. NASLAIN M.G. BADER, J.H. GREENWOOD, D.C. PHILLIPS H. KELLERER, K. FRIEDRICH M. AGNETTI, G. CONNI Th. JOHANNESSON N.SPRECHER P.MElER

B. PIPES, J.C. SEFERIS, S.W.TSAI T. HAYASHI, A. KOBAYASHI, M.UEMERA H.GU

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825

826

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ABISROR ADAM T. AHMED S. ALLAN P. ANDERSEN S.I. AOKI R. ARENDTS F.J. ASLOUN E.M. AVESTON J. BADER M.G. BAILEY J.F. BAPTISTE D. BARBIER B. BATCHEWER D. BATHIAS C. BATOZ J.L. BEVIS M.J. BOMPARD P. BOUlXJ. BOUR J.S. BRANDT J. BRINSON H.F. BROGLY M. BROUGHTON W. BROWN H. BUCHANAN J. BUNSELL A.R. CANTWELL W. CAPRINO G. CARDON A. CARPINTERI A. CASTRILLO M.A. CAVALIER J.C. CHAMBOLLE D. CHANG F.K. CHEN M. CHERFAOUI M. CHITWOOD W.E. CHUBBJ. CLARKE D.A. CLAVEYROLAS G. CLYNE R.T.W. COUTANDB. CURTIS P. CUTOLOD. DALZIEL A. DAVIDSON R. DAVIES G. DAVIES P. DAY D.

301 677 665 375

529,571 521 365 243 713

79,411 87

151 249

759,765 699 733 375 151

257,293 273 365 545 243 741 353

21 111,43]

747 793 545 309 457 99

151 359 87

771 51

451 79

293 205,213,265

233 451 401 587 785 671

419,747 779

De CASTRO P.M.S.T. De MEESTER P. DERBY B. DESARMOT G. DESCHRYVER L. DESTUYNDER P. DI VITA G. DOXSEE L. DRECHSLER K. DUFOUR X. DUGNEO. ELMES D. FAKIROV S. FAVRE J.P. FEEST E.A. FERREIRA A.J.M. FIORE L. FLORIOJ. FORTIER P. FORURIA C. FOURNIER P. FRIEDRICH K. FRIEND C. FROSTIG Y. GABAYSON S.M. GALIOTIS C. GAUDENZI P. GHORBELI. G/BSONG. G/ROT F. GLEZ-VECINO J.A. GOGEVA T. GONZALEZ M.L. GONZALEZ-DIAZ R. GOURSAT P. GRATTIER J. GRENIER P. GRENIER-LOUSTALOT M.F. GUEMES A. GUETTE A. HAINES R. HARRIS B. HENDERSON J.B. HERVE C. HERZOG M. HEYDUCK J. HlEL C.C. HOGGPJ. HOGNAT J. HORSFALL I. HUANG C.L.D. HULL D. IMAZ J. IVENS J. IWATA M. JACOBS O. JACOBS P. JACQUEMET R. JAHANKHAN H.

829

705 465,493,535

199 249,693

111 725 401 603 365 771 129 563 445 693 165 705 609 687 293 389 609

445, 615, 651 227 333 51

765 425 635 803 233 457 445 57

487 93

635 35 35

457,487 121,129

353 677 687 771 43

521 545 809 693 227 441

563, 657, 741 389,643

465 381 615 627 551 765

830

JAMES P.F. JAMET J.F. JANSSENS W. JENNINGS T. JOHNS T. JONES F.R. KAINER K.U. KAMJYA A. KAUSCH H.H. KELLY J.F. K1ESCHKE R. KONIGP. KORINEK Z. KRETSIS G. KREY J. KUMOSA M. KUNZJ. LACOMBE A. LAGRANGE A. LAl D. LANCIN M. LARDEUR P. LE FLOC'HA. LE PETIT CORPS Y. LEEH. LEELJ. LEICY D. LESSARD L. LEVIN K. UAO W.C. UCEAGA J.F. ULHOLTH. UMA A.V. UUQ. LORY P. LUNDBERGR. MACKE T. MADRONERO A. MAIER G. MARCHETTI M. MARKS N. MARQUES A.T. MASARU M.H. MASON J. MASSON J.J. MATTHEWS F.L. MELANITIS N. MENESSIER E. MINES RAW. MOBBS P. MOUNA-COBOS M A MONDRAGON I. MONTE SJ. MORDIKE B.L. MORTON J. MOUUNC. MOURICHOUX H. MURPHY J. NAKANISHI Y.M.

87 5

603 563 353

595.627.665.87.279 171. 221

381 747 347 265

43 287 671 651

657.741 481 99

551 699 273 733 111

185.233 779

315.341 809 359 509 473 389 529 705 72J 721 93

185 193 481 401 587 705 145 213 501 671 765 121 803 395 487

57.643 51

171.221 671 747 257 353 145

NAKANO K. NARDIN M. NASLAlN R. NISHIDA Y. NORTHOLT M. OBERUN A. OSHIMA K. OSMAN M. OUSSET Y. OYTANA C. PAl C.K. PAILLERR. PARTRIDGE I. PAWSON D. PERREUX D. PETERS P.W.M. PICKARD S. PLUVINAGE G. PONTHENIER J.L.

381 243

121.129.157 145

3 15

381 321 725 515 315

121.185 451 279 515 571 199 301

PRENSA MARTINEZ-SANTOS M. PROUHET S.

257 193 129

QUENISSET J.M. REITER H. RICHARD H. ROCHER J.P. ROMER W. ROUCHONJ. ROUGES J.M. ROUUN A. RUBIO A. RUSSELL K. SCHAMM S. SCHMITT C. SCHRODER J. SCHULTE K. SCHULTZ J. SCHWALBE K.H. SCOTT V. SEGAL A. SHARPE D. SIIEINMAN I. SH/AU L.C. SHPlGLER B. SIGALAS I. SILLWOOD J. SIMOES JAO. SMITH R.L. SOMEKH R. STACEY M.H. STEENBAKKER S.L. STEIDLJ. STENZENBERGER H. SUGERMAN G. SUN C.T. TEST F.L. TETI R. TOURATIER M. TOWATA S.l. TRUCHON M.

185.233 677 747 157

43 693 99

419 389.643

779 157 301

171.221 501.615

243 651 139 333 759 333

315.341 177 657 713 705 411 265 65 71

287 43 51

473 687 793 72J 165 771

TRUMPER R. TSAY D.H. TURVEY G. UNDERWOOD F. VALEA A. VALEA PEREZ A. VALENTE S. VALENTIN D. VAN DAELE R. VARCHON D. VERDEAU C. VERPOEST I. VETESNIK P. VIALA J.C. VILATTE M. VINCENT C. VINCENT H. VINCENT L.

139 VIRLOUVET P. 341 WAGNER H.D. 321 WARWICK M. 759 WELLER T. 643 WEVERS M. 57 WHITESIDE J.

309 WITTICH H. 635 WOHLEKE M. 493 WOLSTENHOOME L.C. 515 WORRALL C.M. 431 XIANG Z.

465.493.535.603 YAMADA S.I. 481 YOUNG R. 293 YOUNG R.M.K. 249 YRIEX M.C. 257 7lEGMANN G.

257.293 7lLVAR V. 609

831

451 71

205 333

465.535 779 445 481 4JJ 803 595 165 227 165 635 621 287

833

ll. IE 'ff Q W (()) flU]}) CVD 185 Damage 501.509.521

D N If})IE J! 615.635.699 Degradalion 603 Detection 793 Dielectric Sensor 779 Diffusion 587 Durability 545.551.595

Absorption 627 Dynamic Relaxation 321 Acoustic emission 771 Elastic 151 Adhesion 535 Elastic constant 121 Adhesive 249.785 Elastic modulus 243 Advanced materiol 21 Electrical property 15 Al-Cu motrix 165 Electrochemical 249 Alumina fibre 65.227 Energy 347 Aluminium 151.157.199 Energy absorbing 657 Amine 35 Epoxy 51. 457. 481. Analysis 353 587. 603. 643. Anisotropic 341.441 651.657.677 Aramid Epoxy 603 Epoxy laminate 615 Aramid fibre 3 Failure 321.481.551 Auger 129 699.713 Bismoleimide 481.563 Failure mechanism 309 Boron Nitride 129 Fatigue 145.451.457 Carbon 99.205 501.515.529 Carbon fibre 51.87.243.249. 563.609.677

293.381.411. Fiberglass 51 425.431.465. Fibre bundle 445 481.501.509. Fibre optic 759 563.595.603. Fibre-reiriforced 615. 741. 759. metal microstructure 139 809 Filament winding 401

Carbon fibre/PEEK 747 Finite element 341.353 Carbon-Epoxy 465.643.699 Flexibility 635 Carbon/glass 677 Fracture 111.151.287 Carbon/Kevlar 677 481.571.603. CARP Code 771 635. 651. 705 Ceramic 99.381 Free-edge stress 725 Ceramic composite 5.87 Fretting fatigue 615 Ceramic fibre 15.177 Fretting wear 615 Ceramic motrix 99.121 Geometry 347 CFR 425 GFRP 609.793 CFRP 487.521.571. Glass 713

693.741 Glass ceramic 121 Cluuacterisation 71. 129. 293 Glass epoxy 301.411. 451.

643 515 Characteristic 257.395.481. Glassfibre 347. 375. 603

627. 693 Glass fibre resin 279 CHARPY test 121 Glass-ceramic motrix 121 Compound 389 Glass/carbon hybrid 671 Compression 333. 359. 621. Glass/Polyester 529

677.699 GRP 705.713 Compression fatigue 521 GRP fatigue 457 Corrosion 279 Heat-treatment 227 Crack analysis 309 IIigh performonce 643 Crack growth III High speed Crack simulation 309 photo graphic 803 Cross-ply cracking 571 High temperature 5.99.587 Crystal phase 381 Hot working 233 Crystallinity 809

834

Hybrid 165.677 Microstructure 15.145.233 Hybrid glassfibre 665 287.431 Hygromechanical 603 MLFM 375 Hygrothermol 635 Modal parameter 425 Impregnation 365 Modulus 65.473 Infiltration 139 Moisture 587. 603. 621. Injection 287 627.643 Injection moulding 375 Monofilament 265 Interface 129. 465. 609 Morphology 389

693 Moulding process 375 Interfacial 87.265 Organic motrix 249 Interfacial strength 279 Orthotropic 309 Interferometry 785 PEEK 809 Interlaminar fracture 419. 747 Photographic 803 Interlaminar mode 445 Physicochemical 157.635 Internal strain 765 PMR-15 595 Interphase 129.243 Polybismoleimide 43 Intralaminar 741 Polycarbonate 651 K2ZrF6 157 Polydiacetylene 759 Kevlar 51 Polyester 51.57 Kinetic 35. 57.595. Polyester laminate 551

643 Polyester resin 347 Laminate 315,341.359. Polymer 93.431.445.

375.425.451. 545.621 465.473.487. Polymer material 3 501.563.571. Polymer matrix 79.741 587.595.643. Polymeric 627 665.671.677. Polymeric fibre 71 721. 759. 803 Polypropylenel glass 287

Lithium alumino Polysilazane 93 silicate 121 Polyurethane 51 Magnesium 171.205.221. Porosity 487

293 Post-buckling 333 Mathematical Powder metallurgy 171.177. analysis 353 221 Matrix cracking 493 Powder 445 Mechanical 71. 213. 375. PPS 431

545. 621. 643 Preimpregnated Mechanical behavior 699 polester 57 Mechanical property 15.65.165 Prepolymer 35

193.365.389. Pressure 635 441.465.635 Property 3.705 643 Pyrocarbon 121

Mechanical testing 293 Pyrolysis 93 Mesopore 65 RAMAN 759.765 Metal 151.165 Refractory 99 Metal matrix 177 Rheocasting 233 Metal microstructure 139 Rupture 199 Metallurgical 171 Sandwich 365.509.535 Microcracking 111 733.803 Microdamage 665 Scanning electron Microelectronic microscope 287 dielectric sensor 779 Sheet moulding Microfractographic 301.481 compoung 389 Micromechanic 451.665 Si3N4 145 Microphotographic 71 SiC 93.99.121. Microradiography 375 129. 157. 199 Microscopy 129.287.487. 205.213.221 Microstructural 205.213.287 233.257.265

835

273 Therrrwset resin 43.51 SiC/AI 165 Therrrwsetting 493 Silicate 121 Thermostability 5 Silicon carbide 79.381 Thick laminate 473 Silicon carbide Three-dimensional 473.535 whisker 151 Titanium matrix 185 Silicon nitride 381 Torsion 515 SIMS 273 Transmission Electron SMC 57.389 microscopy 129 Sol-gel 65.87 Trigger mechanism 657 Solid therrrwplastic Twin skinned GRP powder 445 laminate 803 Spectroscopy 129 Ultra-lightweight 395 Sputter deposition 265 Ultrasonic 487.521 Stability 621 Ultrasonic Waveform Static strength 509 detection 793 Strength 151.359.441. Vapour deposition 257

457 Vibration 733 Stress 353. 411.713 Vibratory treatment 177 Surface treatment 279 Video Tensile 705 microphotograpic 71 Tensile fatigue 451 Vinyl ester 51.279 Tensile property 501 Viscoelastic 243 Tension 515.677 Electrochemical 249 Thermal cycling 145 Viscoelasticity 3.545 Thermal degradation 595 Viscosimetric 57 Thermal equilibrium 687 Thermal expansion 121 Thermally 687 Therrrwmechanical 185.545.595 Water 571 Therrrwelectric 563 Whisker 145.151.165.171 Therrrwplastic 375.431.493 205.213.

809 Zinc matrix 193

836

Les actes de ce colloque ont ete imprimes it partir des textes rediges par les auteurs. Les editeurs et les organisateurs ne peuvent iUre tenus pour responsables des opinions exprimees dans ces communications et deserreurs eventuelles qu'elles pourraient contenir.