UCRLJC-12SOS6PREPRINT

Therrno-Mechanical Processing andProperties of a Ductile Iron

CKsynD. R LesuerO. D. Sherby

This paper was prepared for submittal to theThermomechankal Processing and

Mechanical Properties of Hypereutectoid SteelsTlklS Sympc&mIndianapolis, IN

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July 14#1997

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THERMO-MECHANICAL PROCESSING AND PROPERTIES

OF A DUCTILE IRON

C. K. Syn*, D. R. Lesuer* and O. D. Sherby**

*LawrenceLiverrnore national Laboratory, Liverrnore, CA 94551**Dep’t of Materials Science& Eng., Stanford University, Stanford, CA 94305

Abstract

Therrno-mechanicalprocessing of ductile irons is a potentialmethod for enhancing theirmechanicalproperties. A ductilecast ironcontaining3.6% C, 2.6% Si and 0.045% Mg wascontinuouslyhot-and-warmrolledor one-steppress-forgedfroma temperaturein the austeniterange(900°C-1100°C) to a tenqymtwebelowtheAl temperature.Variousamountsof nxluctionwereused(from60% to morethan90%) followedby a shortheattreatmentat 600°C. The heattreatmentleadto a structureof finegraphitein a matrixof ferriteand carbides. The hot-and-warrnworked materialsdevelopeda pearliticmicrostructurewhile the press-forgedmaterialdevelopeda spheroidite-likecarbide microstructurein the matrix. Cementite-denudedferritezonesweredevelopedaroundgraphitestringersinthehot-and-warmworkedmaterials,but suchzones were absent in the press-forged material. Tensile properties including tensile strength andtotal elongation were measured along the direction parallel and transverse to the rolling directionand along the direction transverse to the press-forging direction. The tensile ductility andstrength both increased with a decrease in the amount of hot-and-warm working. The press-forged materials showed higher strength (645 MPa) than the hot-and-warm worked materials(575 MPa) when compared at the same ductility level (22% elongation).

This work was performed under the auspices of the U.S. Department of Energy by the LawrenceLiverrnore National Laboratory under Contract No. W-7405-ENG-48.

Introduction

It is the purpose of this paper to discuss some of the mechanical properties of ductile iron aftervarious thermo-rnechanical processing steps. Although ductile iron is usually used in the as-castcondition, there is now evidence that such a material can have considerably improved roomtemperature properties after various hot-and-warm working steps ‘14].

Ductile ironcan beclassifiedas a hypereutectoidsteelsinceitstypicalcompositionis about3.5%carbonwithabout2.5% silicon. It has many of the characteristicsof ultrahighcarbon steels(UHCSS) in thatwhenheatedhighin the austeniterange, it can have as much as 1.6% C insolution.Thus,thetransformationproductsandmechanicalprocessingstepsstudiedextensivelyin UHCSSU]can be directlyappliedto achievingunusualstructuresand propertiesin ductileirons. It is thepurposeof thispaperto showsomeof the structuresandpropertiesobtainedin aductileiron afterdeformationby hot-and-warmworkingrollingand by press-forging. Someattentionwillbepaidto the formationof ferrite-ffeeregionsduringmechanicalworking,andtorelatingthis structuralfeatureto the diffusionkineticsof carbon and iron in the austeniteandferriteregions.

Jron-Granh ite. Iron-Ir on carbide Phase Diamam

The selection of appropriate them-m-mechanical working steps requires knowledge of the phasesin ductile iron as a function of temperature. Accordingly, Fe-graphite and Fe-FesC phasediagrams for an Fe-2.5% Si-C material are shown in Fig. 1. The phase diagram for Fe-graphite(in dotted lines) is from reference [8], whereas the diagram for Fe-FesC (in solid lines) wasconstructed by interpolation of the diagrams for the Fe-2% Si-C[9]and Fe-3Y0 Si-C[iO]systems.

1 r I.., ..

1200 +. y+L

E ‘i-.”’= ’’’’’:::: A::::?’x:: F---”-”-------”$1.

Y

y + Fe$C

\, -- . . . . . ..-n-..y+ L+ Fe3C

4

fli!5Ld0 1 4 5

~arbon, W&

Fig. 1. Estimated phase diagrams for the Fe-Fe3C (in solid lines) and Fe-graphite (in dottedlines) system containing 2.5% Si.

.@

Several observations that are relevant to the current study can be made from Fig. 1: (1) it definesthe amount of C in solution in austenite as a function of temperature (since it is about the sameline for both Fe-C and Fe-Fe3C); (2) it defines the boundary above which the ferrite (c%)phasedisappears (about 900”C for our 3.5% C ductile iron); (3) it indicates that above 900”C, austeniteis likely in metastable equilibrium with graphite and iron carbide just as, at low temperature,ferrite is in metastable equilibrium with iron carbide and graphite; (4) it predicts the meltingtemperature or start of melting for the ductile iron at about 1120 to 1140”C and indicates theupper limit of metal working; and (5) it suggests that a four phase region likely exists duringmechanical working of the ductile iron from the austenite region to the ferrite region, namelyferrite, austenite, graphite and iron carbide.

Material and Ex~erimental Procedures

~

Ductile iron with a nominal composition of 3.6% C, 2.6% Si, 0.045% Mg and balance Fe wasreceived in the form of rectan&darplateswithdimensionsof 25 mm X-75 mm x 300 mm.Materialwasobtainedin the as-cast and austemperedconditions. Smallblocksof 25 mm x 32mm x 38 mm size were slicedfrom the platesand used as the startingmaterialfor thermo-mechanicalworking.

?roc essinz Proced ures

Sampleswerethermo-mechanicallyprocessedusing rolling and pressing procedures. Details ofthe procedures am described elsewhere[c] but can be summarized as follows: Four differentprocessing procedures, involving either hot and warm working (HWW), or multiple warmworking, were used: (1) Continuous HWW rolling at 1100 to 800”C in seven passes resulting ina true strain of E=l.7; (2) Multiple HWW Rolling to get large strains. The sample wasdeformed to E=l.2 in six passes at 1100 to 800”C, reheated to 1100”C and deformed to e= 1.45in two passes with a total &=2.65; (3) Press-Forging at 1100”C to 8=2.07 in two seconds; (4)Multiple Warm Working Procedure. Samples were heated to 900”C to fully transform ferrite toaustenite containing 1.090 carbon, with graphite and probably some undissolved iron carbide.Four rolling steps, with reheating to 900”C, were used. Each rolling step involved four passeswith e==. 12 per pass, resulting in a total deformation of E=1.80.

The maximum thickness reduction rate (true strain rate) of thickness in each rolling pass wasestimated to be about 0.1 see-i. The forging strain rate was estimated to be about 1 see-l. Alltherm~mechanically worked samples were annealed for 20 minutes at 600”C. The purpose ofthk treatment was to transform any retained austenite to ferrite plus carbide. The presence ofsome retained austenite is expected because of ferrite formation during mechanical working andthe rejection of carbon into the remaining austenite, making for a high carbon austenite, whichremains stable at room temperature.

Tensile Test and MetallomaDhy

Tensile test samples with a 5 mm gage width and 12.5 mm gage length were prepared from therolled or pressed materials by electro-discharge machining (EDM). Samples from the rolledmaterials were machined in both the longitudinal and transverse directions. For the press-forgedmaterials, tensile samples were machined with the tensile axis parallel to the radial direction. AUtensile tests were performed at room temperature. Optical and scanning electron microscopy ofas-received and thermo-mechanically processed and heat treated materials were performed.

Results and Discussion

Microstructure

Typical optical microstructure of the as-east and as-austempered materials in the as-reeeivedcondkion are shown in Figs. 2(a) and (b). Examples of microstnrctures obtained aftercontinuous HWW and press-forging of the m-cast material followed by the short annealingtreatment are shown in Figs. 3(a) and (b), respectively. Meehanieal working (and annealing)

F]g. 2. Optical microstmctures of the (a) as-cast and (b) as-austempered materials in the as-received condition.

leads to the following general rnicrostnrctural features: ( 1) highly elongat&i graphite nodularparticles, increasing in aspeet ratio with increasing mechanical working (Fig. 3(a)) with theexception of the press-forged sample (13g. 3(b)); (2) cementite-denuded ferrite zones are evident,usually located adjacent to the graphite stringers in the HWW treated material, but not in thepress-forged one; (3) the dark appearing regions seen in the micrographs of Figs. 3(a) and (b)(that are not graphite nodules) are regions containing ferrite and cementite as shown in thecorresponding high magnification scanning electron micrographs Figs. 4(a) and (b). The HWWsample shows a microstructure consisted of fine pealite colonies (Fig. 4(a)). The press-forgedmaterial (Fig. 4(b)) shows a spheroid]zed carbide structure expeeted from an uphilltransformation of retained austenite as described elsewhere[c]. The retained austenite resultedfrom the rapid cooling from 1100”C during press forging.

Tensile Properties

The tensile strengths and tensile ductilities of the mechanically processed ductile iron are shownin Fig. 5 plotted as a function of the degree of mechanical work, IZW= -ln (hO/hf), where ho is theinitial thickness and hf is the final thickness. As can be seen, there is a strong effeet of the

Fig. 3. OpticaJ microstructure obtained after (a) continuous HWW and (b) press-forging of tbeas-cast material and the short annealing treatment.

Fig. 4. SEM photomicrographs (a) and (b) of the dark appearing regions in themicrographs of Figs. 3(a) and (b) (that are not graphite nodules), respectively.

processing route and the amount of deformation during processing on the strength and ductility.BotJ the strength and the tensile ductility are seen to decrease with an increase in the amount ofwork. This correlation between strength and ductility is different from that typically observed inductile irons in which the tensile ductility decreases as the strength is increased”]. The basis ofthe unexpected result obtained here can be understood in terms of the materiafs resistance tofailure. If failure is delayed, the strain to fracture will be higher and, therefore, a higher fracturestrength will result. This concept can be tested with the microstructure-and-property resultsobtained here and has been discussed elsewhere[c].

.

TensileStrencXh.

Rolled/Long.

\

.‘.O Pressed

‘.

‘-* .’..

Rolled~r&. = -\.\

‘* -al=(nc QLmli!Y~ 300 –

- 40Rolled/Long. Pressed $?

: Y..; -

--+&-:% - E20 “g

Roiled/Trans.. -.. ~u

Oy ‘ 1 I 1 1 I 1 a i 103

‘w = - & (hO/hJ

Fig, 5. Tensile strength and ductility vs.working strain for ductile iron.

In Fig. 5, the longitudinal samplesshow both a higher stnmgth andhigher ductility than the transversesamples. This result is readilyexplained by the morphology of thegraphite relative to the test direction.In the transverse sample, the cross-sectional area of the graphite normal tothe tensile axis is larger than the cross-sectional area of graphite in thelongitudinal sample. The larger cross-sectional area of graphite in thetransverse sample produces higherlocal stresses in the matrixsurrounding a cracked graphite ribbonthan in the longitudinal sample. Thus,damage development and failure willoccur at lower stress, leadhg to lowerductility and lower strength values inthe transverse sample. Anothervariable to consider is texturedevelopment during rolling. Nostudies on texture, however, havebeen done on the rolled ductile ironand therefore this variable remains tobe explored.

The press-forged ductile iron showsthe highest tensile strength whencompared at the same mechanicalworking strain, see Fig. 5. The tensile

strength and total elongation data from Fig. 5 for the different processing conditions werereplotted in Fig. 6. The press forged material also shows the highest tensile strength whencompared at the same elongation or ductility level as seen in Fig. 6. The higher strength resultsfrom the smaller nodule sizes that are observed in the press-forged material relative to the rolledmaterials. Smaller graphite nodules are obtained because of carbon dissolution at the temperature

(—j~O 5 10 15 20 25 30 35

Elongation, 9’.

Fig. 6. Tensile strength vs. elongation for press-forged ductile iron studied.

of pressing (1100”C) and subsequentrapid pressing. Since the sample wasrapidly cooled during pressing, ahigher concentration of carbon wasretained in austenite during coolingthan in the rolled material. Thedissolved carbon became iron carbkieduring the heat -ting step at 600”Cand dld not increase the volumefraction of graphite. Another reasonfor the higher strength is the absenceof a cementite denuded ferrite zonearound graphite stringem in the press-forged material as shown in Fig. 3(b).The spheroidized carbide structuredeveloped right around the graphitestringers is inherently stronger thanthe cementite-denuded carbon-free

ferrite zones present in the rolledmaterials. 900 1 t 1 I 1 , , I I 1 , 90

/—\lu /

Extrapolation of the data given in Fig./-

% /5 to lower working strains suggests an

\:

● Strength – 6(3~interesting trend in strength and

~- 600 -g

ductility at strains below 1.7 as shown ~ ● .$in Fig. 7. It seems reasonable that the

/-\ =a / G

trend of increasing strength and g~ - /—

1 , ,,rw

3oaductility with decreasing working ~m

. Ductility

strain may reach a maximum and then 1-decrease to the levels of as-caststrength and ductility. Further study is 00 1 2 3°waminted to develop a processing

‘w =-ln(hO/hf)condition which will give an optimalstrength and ductility at a lower strain Fig. 7. Extrapolation of the data given Fig. 5 to

level than used in the present study. low working strain.

Cementite De nuded Zon qe

The absence of cementitedenuded zones in the press-forged material (Fig. 3(b)) and theirpresence in the HWW (Fig. 3(a)) material can be explained qualitatively from the phase diagramshown in Fig. 1. The materials were initially heated to 1100°C and the structure consisted ofaustenite, cementite and graphite, as can be seen from the phase diagram. During the thermo-mechanical processing, as each material cools down from 11OO°C,carbon in the austenite eithermigrates to the nearest graphite nodules or forms more cementite (proeutectoid carbide) particles.However, the large deformation during the thermo-mechanical processing would prevent thecementite formation and even enhance the dissolution of the existing cementite particles byproviding fast diffusion paths resulting from dislocation production. Since graphite isthermodynamicallymore stablethancementite,deformationat elevatedtemperature,and slowcoolingwillalwayspromotedissolutionof cementiteparticles. In thepresentcase of the HWWmaterial,the slow plasticdeformationcombinedwith slow coolinglead to the creationof thedenudedzones aroundgraphitenodules. On the other hand, the fast deformationand rapidcoolingof thepressforgedmaterialpreventedthecementiteparticlesfromdissolving. The rapidcoolingalsopreventedthe austenite-to-ferritetransformationandleadto a largevolumefractionof retainedausteniteor martensitein the as-forgedcondition,whichin turntransformedto finespheroidizedcarbidestructureduringpostheattreatmentas showninFig. 3(b).

The kinetics of denuded zone formation can be understood in terms of the dissolutionmechanisms of carbon and iron during the thermo-mechanicd processing. ‘I’he dissolutionprocess is controlled by the slower of the two diffusing species, iron and carbon. The averagewidth, k, of the denuded zones in Fig. 3(a) was estimated to be about 6 pm and the totaldeformation time, t, during the HWW processing was about 20 seconds. Thus the diffusivity,D, required for rate-controlling species to move a distance of 6 pm was estimated to be about4.5.10-13m2/sec. from the simple relationship, 1=2.(D”t)”2. This calculated value of diffusivitywas compared with the known diffusivity data for carbon and iron in ferrite or austenite for thetemperature range of thermo-mechanical processing used in the present study as shown in Fig. 8.Difisivity data in Fig. 8 were calculated from the pre-exponential factors and activation energiescited in reference 12. A horizontal line representing the calculated dlffusivity 4.5.10-i 3 m2/sec.was drawn in Fig. 8. One can see that the carbon diffusivity is more than an order of magnitudehigher than thk value over the temperature range of the thermo-mechanical processing, and henceit cannot be the rate-controlling mechanism. The carbon, once dissolved, diffuses rapidly to thegraphite boundary. Therefore, since iron moves slower than carbon, it must be the irondiffusivity that controls the dissolution rate of cementite particles. The caicula~d dlfisivity is,

however, much higher than thediffusivity of iron-in-austenite andiron-in-ferrite. Thus, it must be astrain-enhanced iron diffusion that istaking place, either by creation ofnumerous dislocation-short-circuitcontribution or generation ofvacancies[13]. The denuding processmust be occurring in the temperaturerange of 750”C to 950”C where lotsof iron carbide meciDitates will bepresent as can k inkrredphase diagram in Fig. 1.

1.

2.

3.

4.

5.

co nclusions

from the

A ductile cast iron has beenshown to be readily therrno-mechanically worked bycontinuous hot-and-warm rollingand by one-step large strainpress-forging.

In the strain range studied(CW=l.7 to 2.5), tensile strengthand ductility both increase with adecrease in the amount of hot andwarm working.

The press-forged ductile ironshowed higher strength than the

1200C 11OOC 1000c 900C 800C

~’” l“’’I’’’’ 1’’’’ 1’’’’ [’” ‘

104 T

10-” ~T#.1

E

.2%.g , *-13

a=.-0

s

10“’5r

10“’7 11 , # 1,,111111111111111111111 1

6.5104 7.5104 8.5104 9.5104

VT, K’

Fig. 8. Diffusivity of iron and carbon in ferrite andaustenite irons estimated from data in reference 12.

rolled ductfie iron wh~n compared at the same strain.

Cementite denuded ferrite zones around graphite stringers am formed in the continuos hot-and-warm worked condition but not in the press-forged condition.

The formation of denuded zones are rate-controlled by iron diffusion in iron during thethermo-mechanical processing procedures.

Acknowledgment

The authors are gratefully to Dr. P. H. Mani, Technical Director, Wagner Castings Co., Decatur,Illinois, for providing ductile cast iron materials in as<ast and austempered conditions. Theauthors thank Jim Mathews for help in metal working, Bob Accardo for heat treating, JefferyGroth for sample machining, George Yanes and Reynold Lum for tensile testing, and JimFerreira and Ed SMM1Ofor optical and SEM metallography. This work was performed under theauspices of the U.S. Department of Energy by the Lawrence Liverrnore National Laboratoryunder contract No. W-7405-ENG-48.

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References

A. MatSushi@W. Yamamoto,K. HosooandK. Ikaw%Proc. 93rd Meeting of the JapanFoundrymen’s Socie~, 1978, The Japan Foundrymen’s Society, Tokyo, Japan, p. 13 (inJapanese).A. Kagawa, T. Okamoto, H. Matsumoto, and K. Kamei, Journal of the JapanFoundrymen’s Society, 1983, vol. 55 (10), pp. 597-602 (in Japanese).D. J. Moore, T. N. Rouns and K. B. Rundman, J. Heat Treating, 1985, vol. 4(l), pp. 7-24.K. Nagai and K. Kishitake, J. Japan Foundrymen’s Society, 1990, vol. 62(5), pp. 338-343 (in Japanese).S. Yamada and T. Kobayashi, J. Japan Foundrymen’s Society, 1995, vol. 67 (6), pp. 385-390 (in Japanese).C. K. Syn, D. R. Lesuer and O. D. Sherby, Metall. Mater. Trans. A, 1997, vol. 28A (5),pp. 1213-1218.D. R. Lesuer, C. K. Syn, A. Goldberg, J. Wadsworth and O. D. Sherby, JOM, 1993, vol.45 (8), pp. 40-46.D. B. Craig, M. J. Hornung and T. K. McCluhan, “Gray Cast Iron”, from Casting, vol.15, p. 629, Metals Handbook 9th Edition, ASM, 1988.From Fig. 1 of “Basic Metallurgy of Cast Irons”, in Vol. 1, Properties and Seleetion: Ironsand Steels, Metals Handbook 9th Edition, ASM, 1977.0. D. Sherby and T. Oyarna, U.S. Patent No. 4,533,390, Aug. 6, 1985.I. C. H. Hughes, “Ductile Iron”, in Casting, Vol. 15, pp. 647-665, MetaZs Handbook 9thIMition, ASM 1988.Y. Adda and J. Philibert, La Diffusion duns les Solides, vol. 2, Institut National desScienees et Techniques Nuch%ires, SaelaL & Presses Universitaires de France, Paris,France, 1966.M. J. Harrigan and O. D. Sherby, Mater. Sci. Eng., 1971, vol. 7, pp. 177-189.

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