the quasi-static fracture behavior of an

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Composites Science and Technology 53 1995) 307-315

0 1995 Elsevier Science LimitedPrinted in Northern Ireland. All rights reserved

ELSEVIER 0266-3538(95)00057-7 0266-3538/95/ 09.50

THE QUASI STATIC FRACTURE BEHAVIOR OF ANALUMINUM ALL OY METAL MATRI X COMPOSITE

T. S. Srivatsan & A. Prakash

Department of M echanical Engineeri ng, The Uni versit y of Ak ron, Ak ron, Ohi o 44325-3903, U SA

(Received 9 March 1994; revised version rece ived 1 April 1995; accepted 6 April 1995)

AbstractA study has been made t o underst and t hemi crostr ucture, t ensil e deformat ion and fractur e char-

acteri stics of an alumi num all oy 2080, disconti nuously -reinf orced w it h vary ing amounts of sili con carbidepart icl es. Resul t s reveal t he ela sti c modul us andstrength of the metal -mat ri x composit e increase w it h anincrease in r ein forcement cont ent in the metal matri x.The increased strength of t he Al /SiC,, composit e isascri bed t o t he compet i ng and synergisti c i nfl uences ofresidua l str esses generat ed as a result of intrinsicdi fferences in t herma l expansion coefi cients bet w eent he composit e consti t uent s and str engthening fro mconstr ained plastic flow and t riaxial it y in the ducti lemat ri x w hi ch i s due t o t he presence of hard and

ela sti call y defor mi ng reinfo rcements. Fractur e on ami croscopic scale compri sed cracki ng of t he indi vi dualand cl usters of part icul es pr esent in t he mi crostr uctur e.Part icle cracki ng increased w it h reinforcement cont entin the aluminum all oy matri x. Final fracture of thecomposit e resulted from crack propagati on thr ough themat ri x betw een part iculat e clusters. The int ri nsicmechanisms and mi cromechanisms cont ri buti ng tostr ength and governing t he t ensil e fr actur e pr ocess a redi scussed.

to the unreinforced alloy,SO while generally main-taining receptiveness to processing and characteriza-tion techniques used for their conventional unrein-

forced counterparts. From a design perspective theattractiveness of choosing MMCs stems from animprovement in specific modulus, i.e. density-compensated increase in elastic modulus. The moduliobtained are greater than those of typical titaniumalloys and only marginally less than those of moststeels. Associated with an improvement in modulusare concurrent increases in yield and tensile strengthsof up to 60%.- Furthermore, the DRA MMCsbased on particulate reinforcements are attractivebecause they can be made with properties that arenear isotropic in three orthogonal directions or in a

plane. Also, essentially conventional fabricationmethods can be used to produce a wide range ofproduct forms, making them relatively inexpensivecompared to the composites that are reinforced withcontinuous fibers or filaments. The whisker-reinforcedcomposites offer the potential for enhanced propertiesbut suffer from whisker damage and breakage duringsecondary fabrication.4,

Keyw ords: fracture behavior, aluminum alloy, metal-

matrix composite, silicon carbide reinforcement

1 INTRODUCTION

Discontinuously-reinforced aluminum (DRA) alloymetal-matrix composites (MMCs) based on particu-late, whisker or short-fiber reinforcements in a2XxX-, 6XxX- or 7XxX-series aluminum alloymatrix have, in recent years, emerged as attractive andviable commercial materials for the automotive,aerospace and other high-performance markets.

These composites are preferred because they offer anumber of advantages such as a 1.540% increase instrength and a 30-100% increase in stiffness compared

The incorporation of discontinuous particulatereinforcement in aluminum alloy matrices has beenshown in some cases to provide noteworthy attributessuch as high abrasion resistance,lh increased elevated-

temperature strength, improved creep-ruptureproperties,* good micro-creep performance, corro-sion resistance2 and enhanced fatigue-crack initiationresistance compared to the unreinforced matrixalloy.2 Furthermore, the DRA alloy-based MMCs canbe synthesized by using standard ingot metallurgy(IM), powder metallurgy (PM) and mechanicalalloying (MA) processing techniques. Each of thesemethods results in a composite having differentproperties. The PM processing route is generallypreferred since it offers a number of productadvantages. Most importantly, execution of the

process in the solid state minimizes the deleteriousreactions between the metal matrix and the ceramicreinforcement, and enhances the range of potential

307


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308 T. S. Srivatsan, A. Prakash

matrix reinforcement combinations. However, theprimary disadvantage with these materials is that theyfrequently suffer from low tensile ductility (elongationor strain-to-failure), inadequate fracture toughnessand inferior fracture resistance compared to theunreinforced matrix alloy.22~27

The objective of the present experimental study wasto evaluate the tensile deformation and fracturebehavior of ceramic particle-reinforced aluminumalloy metal-matrix composites. The tensile propertiesand fracture behavior were evaluated for two differentvolume fractions of the discontinuous ceramicparticle-reinforcement phase in an aluminum alloymatrix. Emphasis was given to the procurement of thematerial with controlled synthesis (same matrix alloycomposition, aging condition, precipitation charac-teristics and hardness) and reinforcement particlecharacteristics (type, size distribution and shape),while varying only the volume fraction of thediscontinuous particulate reinforcements in order torationalize its influence on tensile properties andfracture behavior.

2 MATERIALS 3 EXPERIMENTALThe DRA MMC materials selected for investigation inthis study were based on a powder metallurgyprocessed Al-Cu-Mg-Zr matrix alloy, designated asX2080. This matrix alloy system was chosen since itprovides excellent combinations of strength anddamage tolerance. The nominal chemical composition

of the matrix alloy is given in Table 1. Thereinforcement was an F-600 grade silicon carbideparticulate (referred to as SIC,), the particles having anominal size of 16 pm. The material was producedand made available by the Aluminum Company ofAmerica (ALCOA). Two different volume fractionsof Sic, (0.5 and 0.20) were chosen. The iron andsilicon elements in the alloy are impurities and arekept to a low level to minimize the formation ofcoarse intermetallic phases. Zirconium is a grainrefiner, combining with aluminum to form the Al,Zrphase which precipitates during ingot preheat and

homogenization treatment. The zirconium-containingparticles (Al,Zr) aid in retarding subgrain boundarymigration and coalescence, controlling grain growthand stabilizing a fine substructure.x Copper andmagnesium are the primary strengthening agents.

3.1 Specimen preparation and mechanical testingBlanks of size 150 mm X 20 mm X 20 mm were cutfrom the as-received composite plates using adiamond-coated saw blade. Tensile test specimenswere precision machined from the blanks using adiamond-tipped cutting tool. The specimens weremachined with the stress axis parallel to the extrusiondirection of the as-received X2080/SiC,-T6 compositeplates. Thus, the gross fracture plane was perpendicu-lar to the extrusion direction in each case. Thecylindrical tensile specimens conformed to standardsspecified in ASTM E-8, with threaded ends and a gagesection which measured 6.25 mm in diameter and 25mm in length. To minimize the effects of surfaceirregularities and finish, final surface preparation wasachieved by mechanically polishing the gage section ofthe test specimens with progressively finer grades ofsilicon carbide impregnated emery paper to remove allcircumferential scratches and surface machiningmarks.

Prealloyed X2080 alloy powder was produced by agas atomization process. The prealloyed atomized

Uniaxial tensile tests were performed on afully-automated. 22 kip closed-loop servohydraulicstructural test machine (Instron) equipped with a10000 kg (98 kN) load cell. The tests were conductedin a room temperature environment (300 K, 55%relative humidity). The composite specimens weredeformed at a constant strain rate of 0.0001 s-. Anaxial 12.5 mm gage-length clip-on extensometer wasattached to the test specimen with rubber bands. Thestress and strain parallel to the load line were

recorded on an X-Y recorder equipped with a penplotter.

Table 1. Chemical composition of matrix AI -Cu-Mg Alloy X2080

Element Cu Mg Sr Al

Wt 3x 1.X 0.2 Balance

alloy powder was screened to 32.5 mesh and thencombined with F-600 particulate silicon carbide usingproprietory dry blending techniques. Blending wasaccomplished to facilitate a homogeneous distributionof the Sic,, both at the macroscopic and microscopiclevels, in the X2080 aluminum alloy powder. Thehomogeneous blend was cold-isostatically compactedto provide a body with 75-80% of the theoreticaldensity, degassed and vacuum hot pressed to producea fully dense billet. The degassing step assists inremoving adsorbed moisture from the carbide particlesurfaces. The billet was then extruded at 441C(825F). Precise details of the primary and secondaryprocessing treatments can be found elsewhere.ls,yHeat treatment of the X208O/SiC, compositesconsisted of solution heat treating the extrusions at499C (930F) for 4 h and cold water quenching.Subsequently, the composites were artificially aged at177C (350F) for 24 h to get the peak-aged (T6)matrix condition. The X208O/SiC, MMCs wereprovided by ALCOA Technical Center in the T6condition.


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The quasi- stati c fracture behavi or of an alumi num all oy 309

3.2 Microstructural evaluation and fracture analysisThe initial microstructure of the as-receivedX208O/SiC, composites was characterized by opticalmicroscopy after standard metallographic preparationtechniques. Fracture surfaces of the fractured tensilespecimens were examined in a scanning electronmicroscope (SEM) to determine the predominantfracture mode and to characterize the fine-scaletopography of the fatigue fracture surface. Samplesfor SEM observation were obtained from thedeformed tensile specimens by sectioning parallel tothe fracture surface.

4 RESULTS AND DISCUSSION

4.1 MicrostructureThe triplanar optical micrographs illustrating themicrostructure of the X208O/SiC,/xxp composites areshown in Figs 1 and 2. The Sic, reinforcement phase,in the aluminum alloy metal matrix, were ofnear-uniform size. However, very few of the particleswere found to be irregularly shaped and these weredispersed randomly through the matrix. Seldom wasan agglomeration of the Sic,, of varying size,observed. An agglomerated site consisted of a fewlarger SIC particles intermingled with smaller, uniformand more regularly shaped particles. The degree ofagglomeration or clustering of the Sic, was found to

Shorttransversedirection ST)

Extrusiondirection LT)

Fig. 1. Triplanar optical micrograph illustrating microstruc-ture of the X208O/SiC/15, composite.

Fig.

Shorttransversedirection ST)

Extrusiondirection LT)

2. Triplanar optical micrograph illustrating microstruc-ture of the X208O/SiC/20, composite.

be largely unaffected by an increase in the particle

reinforcement phase in the X2080 metal matrix (Fig.2). No attempt was made in this study to determinethe particle size distributions for the two MMCmaterials. The matrix consisted of very fine grainswhich could not be clearly resolved in an opticalmicroscope at low magnifications ( < X 1000).

4.2 Tensile behaviorThe ambient temperature tensile properties of theX208O/SiC, MMC for the two different volumefractions of ceramic particle-reinforcement phase aresummarized in Table 2. The results are the meanbased on duplicate tests.

4.2.1 Elastic modul usTest results reveal only a marginal increase in elasticmodulus with an increase in Sic, content in thealuminum alloy metal matrix. The value of elasticmodulus, of the X208O/SiC, composites, was providedas a print-out, as an output of the tensile stressprogram, by the computer control console of theInstron servohydraulic materials test machine. Thisvalue was cross-checked for both accuracy andconsistency by measuring the slope of the initialregion of the stress/strain curve, below the elasticlimit. The elastic modulus of the X2080/SiC/20,-T6


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Table 2. Monotonic properties of X208O/SiC metal-matrix composite

SIC, Youngs modulus, Yield stress. Ultimate tensile Fracture stress, RA Tn(A,,IA,)content. E (GPa) uY (MPa) strength, rrljTS (MPa) ul (MPa)

f Ei Y EEi ) 1 )(vol. )

0 721.5 99 iz - - -01

I::;520

20 105 (15) 430 539

0 Values in parentheses are in units of ksi. Tangency measurements based on extensometer trace. Elongation to failure. Reduction in area. Tensile ductility.

composite is 105 GPa which is:

(i) 50% more than the elastic modulus of thematrix alloy with no Sic, reinforcement (70GPa), i.e the unreinforced matrix; and

(ii) only 6% more than the elastic modulus of theX2080/SiC/15,-T6 composite (99 GPa).

4.2.2 StrengthThe increase in yield strength (which is defined as thestress required at a plastic strain of 0.2%) due toincreased addition of Sic, reinforcements was onlymarginal for the X208O/SiC, MMCs. The matrix alloywith 20 vol.% Sic, had a 7% higher yield strength(430 MPa) than the composite with 15 vol.% Sic,(401 MPa). The ultimate tensile strength of thecomposite is only marginally higher than the yieldstrength indicating that the work hardening rate pastyielding is low. The spread in the strength of thecomposite material is quite typical of that observed forprecipitation-hardened aluminum alloys. The ultimatetensile strength followed the same trend as the yieldstrength of the X208O/SiC, composite. The improve-ment in ultimate tensile strength due to increased Sic,reinforcement in the X2080 metal matrix was only4%.

Whereas the yield strength and ultimate tensilestrength of the X208O/SiC, composites show only amarginal increase with an increase in Sic, reinforce-ment in the ductile aluminum alloy matrix, theductility of the X208O/SiC, composites, as measuredby tensile elongation over 12.7 mm gage length andreduction in area, decreases. The decrease in tensileelongation with an increase in Sic, content from 15 to20 vol.% is as high as 30%. An increase in thediscontinuous Sic, reinforcement phase decreased thereduction in area of the composite microstructure by30%. This observation is consistent with an increase inthe volume fraction of the hard and brittle ceramicparticle (SIC,) reinforcement phase in the soft andductile X2080 matrix.

The manner in which particle size and volumefraction affects the ultimate tensile strength can bebest described in terms of work hardening. Beyond

501 6.91 7.84 8.2532 4.80 5.52 5.1

macroscopic yield the stress/strain curve is wellrepresented by a simple power law. It is expressed as

u = K(E,) (1)where K is the monotonic strength coefficient(intercept at plastic strain E,, = 1) and n is the workhardening or strain hardening exponent slope. Themonotonic stress/strain curves for the two compositesare shown in Fig. 3. Increase in Sic, reinforcementcontent in the X2080 aluminum alloy matrix isobserved to have no influence on the monotonic strainhardening or work hardening exponent. n, of theX208O/SiC, composites.

In X208O/SiC/xxp-T6 composites, with large CTEmismatch strain, the plastic deformation of the ductilealuminum alloy matrix, in the presence of thediscontinuous Sic, reinforcements, is nonuniformprimarily due to the elastically deforming particlesresisting plastic flow of the metal matrix. The plasticdeformation induced dislocations or slip dislocationswould become dominant when the plastic strainexceeds the thermal mismatch strain and the twoeffects would then act in synergism so that they can becombined. The increased strengthening, Au, of thematrix of this Sic,-reinforced aluminum alloy MMCdue to dislocation generation, and assuming thesedislocations are uniformly dispersed in the matrix, canbe estimated using the relationship:

AuY = yGb(p)

0 15 SiCp20 SiCp

100 I I0 01 0 1 1 0 1

(2)

.OStrain ( )

Fig. 3. The monotonic stress/strain curves for theX20XO/SiC, composites.


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The quasi- stati c fracture behavi or of an alum inum all oy 311

where Au? is the increase in yield strength of thecomposite over that of the unreinforced matrix alloy,G is the shear modulus (GPa) of the metal matrix, b isthe Burgers vector, p is the increase in dislocationdensity of the composite matrix over that of theunreinforced matrix density, and (Y s a constant and isequal to 1.25 for aluminum.30 Generation ofdislocations due to CTE mismatch, in discontinuouswhisker-reinforced and aligned continuous fiber-reinforced metal matrices, has been found anddocumented by other investigators.- Taya andMorix4 and more recently Mochida et al. have shownthat the punching of dislocations generated by CTEmismatch strain in a particle-reinforced MMC issufficiently extensive to cover most of the matrixdomain. The dislocations generated by the CTEmismatch strain can at best be considered as anexample of geometrically necessary dislocations.36The geometrically necessary dislocations occur inorder to permit compatible deformation in a systemhaving geometrical constraints such as hard Sicparticles which deform elastically while the surround-ing X2080 aluminum alloy matrix is ductile andundergoes plastic deformation. The geometricallynecessary dislocations become essential when defor-mation occurs without the formation of voids aroundthe brittle ceramic particles. The dislocations resultingfrom slip are a function of intrinsic material propertiesof the system and are not dependent on microstructu-ral constraints. The geometrically necessary disloca-tions contribute to the dislocation density in thecomposite matrix.

The increase in flow stress of the discontinuousparticle-reinforced matrix over the unreinforcedmatrix alloy is proportional to the CTE mismatchstrain if the dislocations generated by the CTEmismatch strain are dominant. The mismatch strain,E,, induced in the particle is given by:2

E, = ((Y, - cu,)AT (3)

where (Y,, and IZY~ re the coefficients of thermalexpansion of the ceramic particle and the matrix,respectively (both the matrix and the particle areassumed to be isotropic in stiffness and CTE), and ATis the net temperature change (TO - Tambirnt) whenthe particle-reinforced metal matrix is quenched froman elevated temperature (TO).

Based on the results obtained in this study, thecontributions to strengthening of the Sic,-reinforcedX2080 MMC arises from the concurrent and mutuallycompetitive influences of several of the followingmechanisms:

(a) The dislocations which are introduced duringprocessing are not completely removed by thesolution heat treatment during the T6 process-

(b)

Cc)

(4

(4

(f)

ing treatment. Consequently, they becometrapped at the Sic particles resulting in localregions of high dislocation density, i.e. thedensity is highest near the reinforcing particleor the reinforcement matrix interfaces.Strengthening due to large differences inthermal coefficients of expansion betweenconstituents of the composite, i.e. aluminumalloy and Sic, resulting in misfit strains due todifferential thermal contraction at the Al/SIC,interfaces. The misfit strain and concomitantmisfit stresses generate dislocations. Theincreased dislocation density generated toaccommodate the misfit strains provides apositive contribution to strengthening of thecomposite matrix.Strengthening arising from constrained plasticflow and triaxiality in the ductile aluminumalloy matrix due to the presence of thediscontinuousments 32.37-3Y

particle (SIC,) reinforce-As a result of the elastic particles

resisting the plastically deforming metal matrix,an average internal stress or back stress, a,,, isgenerated.Contributions arising from competing in-fluences of back stress in the plasticallydeforming composite matrix4 and due toplastic relaxation by the formation of prismaticdislocation loops around the hard and brittlereinforcing particles..A small contribution from dispersion strength-ening caused by the presence of reinforcing Sicparticles in the X2080 aluminum alloy metalmatrix and thus the additional stress requiredfor the slip dislocations to by-pass a particle.Intrinsic differences in texture between theX208O/SiC, composite matrix and the unrein-forced matrix material.j2

4.3 Tensile fracture behaviorThe monotonic fracture surfaces are helpful inelucidating microstructural effects on the ductility andfracture properties of the X2080/SiC,-T6 composites.It is fairly well established that the fracture ofunreinforced alloys is associated with events of voidnucleation and growth, with the nucleation occurringat coarse constituent particles present in themicrostructure.42143 An essential requirement for voidnucleation is the development of a critical normalstress across the particle or the particle/matrixinterface.44 In the metal matrix with no reinforcement,the nucleation of cavities and voids occurs byconcurrent and synergistic influences of:

(a) decohesion at interfaces between the brittleparticle and the ductile matrix; and


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(b) cracking of the hard and brittle inclusions thatdeform elastically.

The X208O/SiC, composites examined exhibitedlimited ductility, on a macroscopic scale, with fractureessentially normal to the tensile stress axis. However,microscopic examination of the fracture surface at

high magnification revealed features reminiscent ofboth locally ductile and brittle mechanisms. Rep-resentative fractographs of the tensile fracture surfaceare shown in Figs 4-7.

On a macroscopic scale the fracture surfaces of thiscomposite were flat, but relatively rough when viewedon a microscopic scale (Fig. 4(a)). The fracturesurfaces revealed fractured particles surrounded byductile regions described as tear ridges. Few tearridges were evident on the fracture surface (Fig. 4(b)).The matrix of the X2080/SiC,-T6 composite wascovered with microvoids of varying size. The voids

were intermingled with isolated regions of dimpledrupture (Fig. 4(a)). The constraints in deformationcaused by the hard, brittle and elastically deformingSIC particles in the adjoining soft and ductilealuminum alloy matrix and the resultant developmentof a triaxial stress state in the matrix, aids in limitingthe flow stress of the composite matrix and favors voidinitiation and growth. As a direct consequence of thedeformation constraints induced by the SIC particulatereinforcements, a higher applied stress is required toinitiate plastic deformation in the matrix. Thistranslates to a higher elastic constant and yield

strength of the X2080/SiC/xxp-T6 composites. Underthe influence of a far-field tensile load the voidsappeared to have undergone limited growth confirm-ing a possible contribution from particle constraint-

Fig. 4. Scanning electron micrographs of the tensile fracturesurface of the IS vol.% Sic, composite.

Fig. 5. Scanning electron micrographs of the tensile fracturesurface of the 20 vol.% Sic, composite.

induced triaxiality on failure of the composite matrix.Particle failure is governed by the conjoint influenceof local plastic constraints, particle size and ag-glomeration. The local plastic constraints are particu-larly important for the larger-sized particles andparticle clusters during composite fracture.44.4 Exami-nation of the tensile fracture surfaces revealed damageassociated with fracture to be highly localized at the

discontinuous Sic, reinforcement with little evidenceof void formation away from the fractured SICparticle. Fracture of the Sic, was greater in regions ofparticle clustering due to enhanced local stressesresulting from restriction of plastic deformation. Theintrinsic brittleness of the reinforcing Sic, and thepropensity for it to fracture due to localizeddeformation results in particle cracking being thedominant damage mode. The higher yield strength

Fig. 6. Scanning electron micrographs of the tensile fracturesurface showing fine microvoids and isolated shallow

dimples.


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The quasi- stati c fracture behavi or of an alumi num all oy 313

PI Silicon carbide

Fig 7 (a) Scanning electron micrograph of the tensilefracture surface showing cracked SIC particles. (b)Schematic representation of the fracture mode for

Sic-reinforced X2080 aluminum alloy matrix.

coupled with concurrent damage to the compositemicrostructure from the conjoint influence of particlecracking and decohesion results in a decrease inductility. Furthermore, assuming that thematrix/particle interfaces are strong the triaxialstresses generated during far-field tensile loadingfavors limited growth of the microvoids in the matrixof the composite. The limited growth of voids duringfar-field tensile loading and lack of their coalescenceas a dominant fracture mode for the X208O/SiC/xxp-T6 composites clearly indicates that the deformationproperties of the X2080 aluminum alloy aresignificantly altered by the presence of SIC,. Also, thepresence of the Sic, raises the hydrostatic componentof stress and, hence, their distribution is an importantfactor governing fracture of the composite.4h47

Fracture of the brittle Sic particles and concurrentfailure of the surrounding matrix results in theformation of voids. Very few of the fine microvoidscoalesce and the halves of these voids are the shallowdimples observed on the fracture surface (Fig. 6). Thelack of formation of ductile dimples as a dominantfracture mode is attributed to the constraints onplastic flow in the composite matrix caused by thepresence of the discontinuous Sic, reinforcement andnot to the limited *ductility of the aluminum alloy perse. The constraints in plastic flow favor the formation

of fine tear ridges between the Sic particles. With anincrease in Sic, reinforcement content, fracture wasfound to be dominated by cracking of the SIC particleson account of their intrinsic brittleness (Fig. 7(a)).With an increase in strain, the larger-sized particlesfracture first, followed by fracture of the smaller-sized

particles. In regions of particleclustering or

agglomeration, the short interparticle distance facilit-ates linkage between neighboring voids and cracks asa direct result of decreased propagation distancesbetween the cracked particles. Based on an observa-tion of the fracture surface, it is seen that the fractureplane of cracked particles is perpendicular to theloading axis, suggesting the importance of the tensilestress in inducing particle fracture (Fig. 7(b)). Theearly cracking of the SIC particles is largelyresponsible for the lower tensile ductility of theX2080/SiC/xxp-T6 composites.

The overall damage resulting from uniaxialstraining of the X2080/SiC,-T6 MMC is due to theconjoint action of two mechanisms:

(4

@)

damage associated with the discontinuous Sic,reinforcement, such as particle cracking anddecohesion at the particle/matrix interfaces;andlattice damage such as dislocations and pointdefects, coupled with residual stress effectsassociated with the presence of discontinuousSic particles.45

Damage associated with particle fracture orcracking is well reflected by changes in elasticmodulus. Changes in residual stress levels duringstraining also result in a concomitant change inmodulus. The degree of damage, D , can be expressedas:4x

D = [l - (EinstantlEinitiai)] (4)

where Einitial is the initial modulus and Einstant is themodulus at any instantaneous value of strain. Thecontribution of cracked particles to the Youngsmodulus of a particle-reinforced MMC can beanalyzed by considering the cracked particles as

penny-shaped cracks and the shattered particles ascomplete voids.3For the Sic particles to fracture completely. they

must be loaded to their fracture stress. This isachieved globally by the tensile stress and locally byshear loading through the interface. The extent ofparticle loading by the shear mechanism is dependenton the aspect ratio of the reinforcing Sic,. For thecase of symmetrically packed particles in a metalmatrix, the aspect ratio, S,, for maximum loading is:4

S, = USC/ Ti (5)where uslc is the strength of the particle and ri is theinterfacial shear strength. The strength of monolithicSIC is about 2000 MPa and assuming that r, = a,/2,


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314 T. S. Sriv at san. A. Prak ash

where (TM is the maximum stress achieved in thealuminum alloy matrix (about 500 MPa). the criticalaspect ratio is 8. The shear mechanism approachignores any end-loading effects which would exertadditional stress on the SIC particles. A carefulobservation of the tensile fracture surface revealedless than 50% of the Sic particles had fractured. Thisindicates that not all of the SIC particles were loadedto their fracture stress suggesting the near-uniformdistribution of the SIC particles in the metal matrix.This situation is further complicated by:

the SIC particle clusters, and is in the form of crackedparticles, tear ridges and voids which have formedaround the cracked particles. The fracture initiates byparticle cracking coupled with decohesion of thematrix surrounding and between the particles. Finalfracture is achieved by fracture through the matrixbetween the particle clusters (Fig. 7(b)). Few voidsgenerated by particle cracking did not growextensively in the tensile stress direction, which isgenerally the case in ductile fracture of unreinforcedaluminum alloys (Fig. 8).50~s2 The lack of extensivevoid growth in this Sic particle-reinforced X2080metal matrix also suggests that the fracture strain iscritically controlled by both the void nucleation strainand linkage strain.

(a>

(b)

The mismatch strain and concomitant internalstress in the composite matrix due todifferences in thermal expansion coefficient(6:l) between the Sic particles and thealuminum alloy. Assuming spherical particles,the mismatch strain, E,,, that will be induced inthe particles will be given by eqn (3). For

particle fracture to occur, the applied far-fieldtensile stress will have to overcome anyinternal stress present in the particle.Loading the particles through the misfit straingenerated during plastic flow as a result of thedifference in elastic modulus between the hardand brittle SIC particles and the soft andductile X2080 matrix.

The radial stress, (T,, at the interface of an elasticinclusion embedded in a plastic matrix was found tobe tensile in nature, and is given by:44%y

where v(E~) is the matrix yield stress at the plasticstrain adjacent to the SIC particle, and a=, is thetensile hydrostatic stress developed in the immediatevicinity of the reinforcing particle. When the radialstress generated at the particle/matrix interfaceexceeds the strength of the SIC particle, cracking ispromoted.

For this SIC particle-reinforced X2080 aluminumalloy matrix the majority of damage is associated with

Fig. 8. Scanning electron micrograph of the tensile fracturesurface of an unreinforced Al-Cu-Mg alloy.

5 CONCLUSIONS

Based on the results of this investigation on effect ofSic particle-reinforcement content on tensile behaviorof X2080 aluminum alloy composite, the followingconclusions can be drawn:

4.

5

The as-received microstructure of theX208O/SiC, MMC revealed a near-uniformdistribution of the SIC particles in the threeorthogonal directions of the extruded compositeplate. Seldom was an agglomeration of the SICparticles observed.Increase in the amount of discontinuous Sic,reinforcement in the X2080 matrix increased theelastic modulus and the strength of thecomposite and degraded ductility.The increased strength of the discontinuously-reinforced X208O/SiC, composites is rational-ized in terms of mechanisms based on anincrease in dislocation density in the matrix dueto the presence of the discontinuous SICparticle-reinforcement phase misfit strains andresultant misfit stresses due to differentialthermal contraction at the Al/SIC, interfaces,and constrained plastic flow due to the presenceof the discontinuous SIC particle reinforcementsin the aluminum alloy metal matrix.The presence of the hard and brittle Sicparticles in the soft and ductile metal matrixcaused fine microcracks to initiate at low valuesof applied stress. Fractography revealed limitedductility on a macroscopic scale, but micro-scopically features were reminiscent of locallyductile and brittle mechanisms.Fracture of the matrix between particle clusters.coupled with particle cracking and decohesionof the matrix surrounding the particles allowsthe microcracks to grow rapidly and link byfracture through the ductile metal matrixresulting in macroscopic failure and resultantlow tensile ductility.


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the quasi-static fracture behavior of an

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