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b IA-UR-95-116f)
We:
Author(s):
Subnitt6d 10
LosAlamosNAII[)N AI lAllollAlo FiY”
STRESS RELAXATION IN DISCONTINUOUSLY REINFORCEDCOMPOSITES
RECEIVEDHAY081995
0S11N. shi, ~. .1. Arscnault
ASME (:olltl’~1’11(’1” , “Ml l.ri~m(’rll;lll lrs ;III(I (:ilIIst I tilt lvL’
-.. _. ......_.... .....-,. - .,. ,..?
10RAlmon NmlmmlI dkilnhvy, nn 1111~nmtwtlm:tklrtm~udIIIMnNtIIIIIIYUIIIIIIIIYW,mOPIIIIWI IIV UI* IIIIIVIWW III (MWIIIIII hw III@~I !; lWWiIIWIIl 01 I IIwINIllhlal Cllllllnl:lw /40% I N(1 M I!y MlI qlldlkxl 111Illlnndldu, ltlupu:!hllnl Inl.lkulwnnM 11101I !+ ( h)vwllmmll Inlnolmii ll~mdllnlvw nlpdly fl~ll lltlllln~ l!)
IMIIIIlnh 01 tWIIMtml Ihs PIlhildhul10111{)1lhlnLllll!lh!hll!, III I(I nlk)wI)lhm 1[)(k) no,tot II % I klVnlllllmlll ImIPIln@m 11111I m AlaItIImt4mIIvId I nlh)tllluwlqlldn ltlnt Iha @lhh ktdl!y llIIh.Wlldn❑ mWOlh@MIIWj lntdmIhu ❑llqnl:n~ (11III* I I !; IIHIUIIIIIW IIf I llaloy
I aIIIIINatn.lm11’I
STRESS RELAXATION IN DISCONTINUOUSLY REINFORCED
(;OMPOSITES1
N. Shl
MS H805, LANSCE
Los Alamos National Lab.
Los Alamos, NM 87545
R. J. Arsonautt
Dept. of Matwials & Nuclear Engineering
University of MarylandCollegIa Park, MD 20742-2115
AOSTRACf
It IISS been obscnmd Ibat in dhmonthwously-reinforced At20J1’liAl composites thst as the
rcinkmemcnl size increases lhe m’ersge density O( dislocmimrs gmermed from the relmxmion
of Ihc lhcrrtml slrcsses increases, and Ibc corresponding Ihmrrd miidual rdrcsscs slighlly
decrease. Similm cb-nges msull wbcn Ibe rcinl’orccmcnl morpbtdogy cbangc!i from spheres
to short fi~rs 10 continuous fllamcnls. The cburgcs of dislocmlorr i’cnsily and Ihcrrrml residual
sIrc.sscs wilb respect to pmticle si?c me In contrast 10 Ihosc r,bscrvcd in Ibc SiC/AJ
countcqml. A previously developed slmplc model u.wd to explain Ibe SiC/Al IIaIm, wbicb WM
based on prismatic dislcatlon punching, suggcslcd that the dcnslly of the misfil disloaions
dccrcmtcs when tbc reinforcement size inuca.ws. In ibis Investigmlon, a simplr model is
proposed to cxpl~in tbr mnomdy in tbc develnpmenl of Ihcmml rcsidunl SUCS.USand Ihc
generation of misfit dlslocmlons M a thncllorr O( the ptrtlclc Azc and shnpc in A120#JiA.1
composites. As a result of n lack 0[ sufficient indc~ndenl.xllp. syslems in low symmclry
mstcrimls such ts NIAL plMIlc relwmdon of the thermal strc.ssceis ,sevcrcly constrained as
compared to fcc Al. As such, plm.k rclmxadwr rcquirm collnhorativc sllps in mnmggrcg-tc of
grains. This only rrccurs when the Icngtb wale ot’ the varying midll Ihrnntl stress flcld is
much larger Ihmi Ihc avcrngc @r size. Thw Is, IIJCmcrh~nlsm of plasIic rclnxtllorr lwromes
o~ratlvc whrn Ihc rcinfiwccmcnt size Incrca.ww
1 This resemch was suppnrted in parI hy the Otl’kc of Naval Rcscmh undm granl NWlt)14-
91-J- 1353. N, Shi would Ilkc 10 arkuowirdgc the support irmu Ihc U.S. Deparimt!ru ot’
Energy.
MASTER
lNTRODU~lON
It basbcen proposed tha[dislocaliona arc gcnerwed by relaxation of the thermal sIrcsses.
These [hcmnal strcs,scs arc developed during the cooling of a composite with a reinforcement
and matrix which have different cocfficicn[s of expansion (Arscnault, 1984. This concept has
been used in investigating several composile systems such as SiC/At, Si/Al, AJzOfliA.l and
Tiq/NiAl (Amcnault, lWI, Arscrmult and Fisher, 1983, Vogclsang, et al,, 1986)
several models have been developed [o predict Ihc density of dislocations generated and the
cffcctivc plastic strain (Amcnauh and Shi, 1986, Shi et al., 1992) due to the rcltxation of the
thermal sircascs. The prismatic punching model developed by Arscnault and Shi (1986) is
capabk of predicting how the dislocation density would change in a homogeneous matrix with
inctcasing the volume frac[ion and psrticlc six when the [hcnnal misfit arc compk[cly
released through dislocation generation, A%the volume fraclion increases the density of
dislocation increases, and as [he pafiiclc size increases at a constant particle volume fraction
the dislocation density decreases, These rcsulIs come from the fact thal the dislocation density
arc linearly related to Ihc total particle surface area. Experimentally it has hccn shown that in
Ihc case of SiC/Al [he changes in dislocation density wi[h particle size follow Ihc prcdic[ion
of ArscnaulbShi model (Arscnaulf and Shi, 1986) as shown in Fig.1.
Io’4 m.2
MML 200
g2 -
.g I20 V% SiCp/l 100 Al
8
51.5 ,., ,. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
‘a
8
aal . . ...,,,,,, . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
&
%
2,= 0,5 .,, ,. ... . . . . . . . . . . . . . . . . .,, .,,,. ., .,..,, ,
i ~
b=
O*1 I 1 1 1
0 50 lrxl 159 200 2!KI
Particle size (D) ~m
Fig, I The Increase In average matrix dlslcrcatlondensity In 20 V% SIC/Al compoallein the as annealed (12 hrrr at 530”C), furnace cooled (12 hrs) condltlon as afunction of particle SIZO, The Increase In dislocation density (Ap) IS equat to
bm~sim - I)mablx, The malrlx Is annealed under the rtame rmrrdltlormns the
composltsm
.
t
In polycrystalline low symmetry materials such as NiAl, plaslic I1OWin a paflicular grain
is difficult due to the lack of a sufficient number of indcpcndcnt slip systems and [he
constraints by the neighboring grains. Therefore, in inttnnctallic matrix cumpositcs generation
of Ihcrmai misfit dislocations, may be affected by the slip constraints. In contrast 10
aluminum-based composites where a high mobile dislocation dcnsily is bcnelicial to [he
strengthening, a high dislocation density in irrlcrrnctallicbasedcompositesmay comributc 10Ihc improvement of ductility.
In Ibis investigation, thermal residualstresses(TRS) and the matrix dislocation density in
AfzOJ-reinforced NiAl were dclcnrrincd hy neutron diffraction and Transmissioncleclron
rnlcrmwopc(TEM), rcspcctivcly.The invcstiga[irrnwas conductedto uctcrminc the influence
of parliclc size and shape on Ihc plastic relaxation of Ihc thermal mistit stresses in [Lis
composite. The results were compmcd with those of SiC!Al composites. A complementary
FEM analysiswas conducted,and a simple model was cOnStNctcd 10explain the expctimcntal
results.
EXPERIMENTAL PROCEDURE
Composites of 20 V% AlzO, with five diffcrerr[ ,emforcement sizes and shapes were
produccd(tiqui-~ xetf parlic!cs with rliamctcm of 5, 75, 355 /(m, shorl fibers wi[h diamclcr of
lf)~m and average aspect ratio of 10, and coruinrrous filament Wilh dinmcter of 144 jm). Thecontinuous filament ~:OJ/Ni~ composites were prodtrccd by a powder cloth Icrhniquc,
whcmaa, the cqui-axed parliclcs were produced by mixing powders of AfJ03 and NiAt of the
same size, and then hot prc.aaing nt 169J K for 4 hrs al a pressure of 25 MPn. The short fiber
composite was produced by mixing 75 pm NIAI powder wilh the Al@3 shofi ti~m. followed
by the same hot prcaain~ procedure. After processing, all composites were anrrcalcd al 140t1°C
for 1 to 4 hours followed by [umacc-cool. Since Ihc compmilts (particle and shofi.fitur) were
producedby a similar PM pmccrfurcs, the as-proccs.wd microstructure in the mawix should
IM simil~r, i.e,, similar grain size, etc.
Ilc Iransmimiorrclccmonmicroscopywas pcrfonncd with a 200 KV ml 1 MV TEM. Tfrcdctalls of foil preparation nnd dnta analysis arc given clsewhcrc (Wsn~ ct al.]. The
nlcnsurcments rrf the TRS in Ihe matrix were perfonncd hy ncrrn’ondiffraction in Univcrsily
of Mlssourl Rcscnrch Remww with a munoclrromatic wave length O( 1,2 A, nnd2?0 NiAl
f3mgg peak was used to dctcmune matrix Iauice swains. The prncmlurm arc dcwribcd f’urthcr
by Sbl c! al. [SfIl et al,, 1Y93, Smith cr d., IW2)
In the FEM Invcstlgalion the A13AQUS FEM cofle WM USC4 mI~ nlr~frc~ w~wntinll IWO
dlmen~irmal axlsymrnctrlc unit cells containing spherical, short IIIIJ {wnlinuous tmvlindricnl
inclusion were made. The prmwturc is dcwrilrcd In more drtrnil hy Shi cl al. ( 1W2),
reinforcement increases, Ihe dislocation density increases (Fig.2b). The changes in dislocation
derrsi~ as function of the reinforcement are [be opposilt to ihose ob[ained for SiC/N
composites (Fig.1), in which [he diskwa[ion dcnsi[y decreases as [he SiC size increases.
MML370
~ 10’4 E
$...........
. . . . . . . . . . ,,, .
i
. . . . . . . . . . . . . .
. . ...1.... ., .,,,..
. . . . . . . . . . . . . . . . . . . . . . . . .
:10” ~,- . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . . . . . . . . . . . . . . . . . . . . . . .
Flg2a
. . . . . . . . . . . . . . . . . . . . .,,
,010 ~sphere fiber filament
The Increases In average matrix dislocation density In 20 V’% A120JNIAI
compoattas wtth particle shape (the sphere repremanta 355 ~m near. equlaxed
padlcles). The Increase In dislocation density (AII) Is equal to pmmPd,O -
P~.WIX,The matrix I@annealed under ttle same conditions RS the com~oslte
IJIZXJAIMEK
I%is report was prepared as an account O( work qmnsored by an agency of the United SmlCS
Govcrnmen?. Neither the Unid Stales Government nor any agency [hereof, nor any of their
employees, makes any warranty, exprw or implied, or assumes any Iegtil Iiabilily or respnsi-
ndily for Ihe accuracy, complelcne~, or usefulness of any information, appuralus, product, or
process disclosed, w represents that i[s usc would no[ inrringe privately ownd rights, Refer-ence herein 10 any specific commercial pr~ucl, prm~~, Ur ~miw by lr~dc name, lr~demark,manufacturer, or otherwise does not ncccuarily consl[tulc or imply its endorsement, rccom-
mendalion, or favorin~ by the United States Government or any agency [hereof, The views
&nd opinions of aulhors eaprd br~in dr) not n~~arllY SlaIe Or reflWI [hM of theUnited Slata Govcrnmen[ or any agency lherq,r.
E
c.-
Flg.2b
10’5
1410
10’3
10’2
10”
do
MML41O
I I I
I I II I II I II I II I II I I1 I II I IL J‘---- ----- ----- ---- ----- ----- .-L
I I II I II I II I I
!:I I
I I 1
I I
I I II I
[
----
ii
-- h------- 4 --------+ -----I I 1I II I1 I II I I
II II I II I I
F------ -r ------ -q ----- --- p---- ---
I I I
II I II I II I II I II I II I II I II I IL 4----- ---- ----- ---- ---._- L-------I I II I II I II I II I II I II I II I I
c) 100 200 300 400
Particle Size ( D ) w
Increase In average dislocation denelty In 20V% A120JNIAI composite with
particle elze,
TABLE 1
Thermal Resldr.mlStress In tho NIAI Matrix
Tensile .Wcss (MPa % 20 MPa)
A1203 Reinforamen[Axial Transveme
75* parricle
355 pr parlicle
short tllxr
cxmfinuoua tilament
144
105
280
mmprcwiveII (Saigal and Kupperman, 1901) I -“ 100
171
189
383
‘J8
<
Tsblc 1 Iisls [he Ihennal residual slrcsscs measured by neulron di[fracliort in [kc
composites. The axial and the transvc= tlireclions are defined as Ihc directions pnrallcl and
pcqmdicular to the hot pressingdircclion during processing.The axial l_RS dccrcascsas the
particle sim increaseswhile no clear mend ctin k de[cctcd in Ihc transverse direcliwt (changesarc witiin experimental cmor). Wi[h changes in reinforcement shape Table 1 indicates that
Ihcrc ia a significant mhrciion in [he matrix TRS as the morphology of the reinforcement
changes from Ilbrs 10 equi-axed particles. However, the matrix TRS is leas in Ihe composite
reinforced by conlinuons filaments.
We could not obtain sensible sims values for [he 5#m A1203 compcraile due 10 iron
comaminalion of [he NiAl, which had been ball milled 10 obtain 5 ~~mpwdcr. The TRS is
an.ktropic and the stress values from lransvc~ ditcclion is Iargcr in all cases. While # more
delailed invesliga[lon is king undctiakcn, we believe the following factors may have
contributed to the aniscmopy in TRS: (a) Ihe shorl fibers arc seen as a plnnnr anay
perpendicular 10 the compact ttircctirms. (b) for larger patliclc size (355 A), Iherc was also a
large dcgrcc of reinforcement cluslcriug along Ihe pcrpcndimtlar directions. Anisotropic elaalic
interaction Mwecrt grains In a IcxIutcd mamix CWIInl.somrrmibule 10 the ardsotropy 0[ the
avcrxge TRS,
Figure 3 shows Ihe changes In svcragc maidx effccllvc plastic strain wl’h reinforcement
morphology as deletmlncd by an FEM analysis. There Is about a factor of two incrca.w in
efkctivc plsstlc strain from spherical 10 continuous Illamenl rcinfo~cmcnt. If the cffcuivc
plsstlc slraht scalcn wilh rli!docaliongcncralicm,this resultis conlislcnt wilh lhe TEM rcriulL%.
DISCUSSION
From an initlsl conaidcrstlon of the tlala it is obvinus thal Iherc is a large diffcrcncc hctwccn
SiC/Al and 41zO@lAl comprmilcs, cspcclally in Ibc gcncrallon of dislocalimts. II Is flrsl
necessary 10 consider If Ihe Ihcmtal !urcsacs arc Mflcknl :0 prrxtuc? strcwcs that arc ahovr
the yicltl sltcsa of Ihc nm[rix (NIAI). An shown in Fig. 4 [he calculated Ihrnnsl SWCM(without
plastic rcltxalion) al all Icmpcraturea is Iargc: than [hc ylcld slrc~ !)( prdycryslalllnc NiA1.
Tbctcforc, dlsltwatlon gctiraliou ml mollon sbnuld owur.
n
m‘o
w
c.-td
8
a>---
~
t?
6
5
4
3
2
●●
●’8#*
●
4°●’
●’●*
●⑤●⑤
●’●
8°●
●°●°
●’●
●e
sphere fiber filament
Flg,3 Tho FEM-pradlcted averag~ dfactlw plasflc strain for different reinforcement
b
1,200
1,000
800
400
200
0
Cal, curve of TRS
of Al ~0 ~ /NiA1.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
.----, Exp. yield stress
of NiAl
●. ---
---~●
●●%
. . . ... . .. . . . . . . . . . . .. . . . .. .
b98\9
‘a●O
200 400 600 800 1,000 1,200 1,400 1,600 1,800
TEMPERATURE K
The experimentally meaeured yield strew of polycrystalllrw NIAJas a functionof feet temperature. The predicted matrix thermal stress without plastlc
relaxation,
MML411
particle
“?)
affected /
zone grains
~
particle
.—
(a)
(b)
Fig.5 Mismatch-affected zone In which high thermal stresses develop, (a)the
mismatch affected zone Is much larger than matrix gralrw (large particles); (b)
the mismatch affected zone Is much smaller than matrix grains (small
particles),
The mosl distinct difference in SiC/A.l and A120Y’NiAl are the changes in Ihe mismatchdislocation density with [he partic!e size as can be wen by comparing Figs. 1 and 3. The
increase in the dislocation density in lhe NiAl ma[rix with increasing pafiicle size cannot beexplained based on relaxationof thermal mismatchin a homogeneous medium. To understandthis anomalous trend in the change of dislocation density i[ is necessary to consider the slipcharacteristics in the two matrices.Al hasa fcc crystal structurewith 12different slip sys[emsof which 5 are independent slip systems, this allows an individual grain to slip independently.
In NiAl where the predominate slip system is c1OO>{ 110} there are only three independentslip syslems (Groves and Kelly, 1963) wi[h [he hard orienla[ion along <100> and the softorientations <110> and <111>. The difference in the cri[ical resolved shear stress between the“hard” and “soft” orienla[ions is atmut a factor of fourteen (Noebe et al., 1993). According toVon Mises (1928) however, slip in individual grains of a polycrystalline ma[erial withoutsacrificing intergranual deforrnaiion compatibility requires five independent slip syslems.Therefore, slip wi[hin a grain without crea[ing disconlinui[y in the NiAl matrix can occur only
when collaborative slip from neighlmring grains is activated. Such a de fomlation mode canbe facilitated when the misfitting themlal slress field e.lcompasses the pZkIeS over asignificant number of matrix grains. Figure 5 shows two extreme cases in which thereinforcement size is much larger (Tig. 5a) and smaller (Fig. 5b) than [he surrounding matrixgrain size. Due to the thermal mismatch stresses a misrna[ch-affected zone arises around each
particle within which a high thermal mistil s[ress develops. When the grain size is muchsmaller relative to the mismalch affected zone, i.e. large panicle size (Fig, 5a), criticalresolved shear strtss is exceeded in many grains. In some grains [he crystallographicorientations are aligned favorably for collaborative multigrain slip without tlcs[roying intergraincompatibility. Wirh a larger grain size relative to the mismatch affected zone, i.e. smallreinforcement size, grain alignment favorable for collaborative slip is statistically unlikely,
Therefore, collatmmtivc slip in an aggregate of grains is possible only when the size of thernismatch-affcctcd zone is much Iargcr [ban [hc average grain size (Fig. 5a), i.e., plasticrelaxation is more Iikcly with a larger avcmgc parliclc size. FEM modeling using a
polycrystalline aggregate model is needed ro quamilatively understand this phenomenon.
CONCLUSIONThc following conclusion can be drawn from [he experiment da[a and [hc modeling.
_ The dislocation density, due to the relaxation of the thermal srrcsses, increases as the A1203size increases in Al,03/NiAl composites. In contrasl, in [he case of SiC/A.l composites as SiCparticles size increa-ses the dislocation density decreases.● The distinct contrast in the dislocation density/particle rcla[ionship for [he two composi[cscan be explained by a simple model based on [he requirement of collabomlive slip Ixrwcengrains in polycrystalline NiA.1.● FE&f results have has shown tha[ shape charges from a sphere 10 cominuous filament
induce afmN a factor of P.vo increase in the effeclivc plas[ic s[rain, which has the same trendwith the changes in dislocation densi[y obtain from TEM.
● Neumon dlffrnction results indicate thrd the mutrix Ihemlnl residual stress in AIJO:/NiAlcomposites decreases as the purtlclc size increases.
REFERENCESArsenault, RJ. and Fisher, R. M., 1983, “Microstructure of Fiber and Particulate Metal
Mamix Composites”, Scripts Metal., Vol. 17, pp. 67-71.
Amcnault, R.J., 1984, ‘The !hengthenirrg of f3161 Aluminum by Fiber and Plmelet SiliconCarbide”, Murer. Sci Eng., Vol. 64, pp. 171-181.
&senault, RJ. and Shi, N., 1986, “Dislocation Genera[ion Due 10 Differences in
Coeffrcienrs of Thermal Expansion”, Mater. Sci. Eng., Vol. 81, pp. 175-187.
Arsenault, RJ., Wang L. and Feng, C. R., 1991, “S[reng[hening of Composites Due ToMicrostrucrural Changes in the MatriY”, itcm Me[al/., Vol., 39, pp.47-57.
Groves, G.W. and Kelly, A., 1%3, Phil Msg., Vol. 8, p. 877.Mises R. vorr, 1928, Z. angew. MatIi. Mech., Vol. 8, p.161.
Noe&, R. D., Bowman, R.R. and Nathal, M. V., 1993, “Physical and Mechanical Propertiesof the B2 Compound NiAl”, Inter. MIiI. Rev., Vol. 38, pp. 193-232.
Saigal, A. and Kupperman, D.S., 1991, Residual Themml Strains and Stresses in Nickel
AJuminide Matrix Composites”,Scripts MetalI., Vol. 25, pp. 2547-2552.
Shi, B., Wilner, B. and ilrsenaul~ RJ., 1992, “A FEM SIudy of [he Plaslic Defonnalion
Process of Whisker Reinforecd SiC/Al Composites”,Acts A4em/1.,Vol. 40, pp. 2641-2854.Shi, N., kserraul~ R,J., Krawilz, A.D. and Smi[h, L.F., 1993, “Deformation induced
Residual Stress Changes in SiC Whisker Reinforced 6061 Al Composites”, Mera//. Trans., Vol.
24Aj pp. 187-196.Smi[h, L. F., Kawitz, A, D., Clarke, P., Saimo[o, S., Soi, N. and Arserrault, RJ., 1992,
“Residual Sresses in Discontinuous Me[al Ma[rix Cbmposi[es”, A4uI.ScL& Eng,, Vol. A159,
pp. L13-L15.Vogelsang, M,, hrurult, RJ. and Fisher, R. M,, 1986, “h in situ HVEM study of
dislocation generation at al/sic imerfaces in melal matrix composites”, MefaL Tram, 17A379-389.
Wang, L., Bowman, R,R. and Amenault, RJ., “Dislocation in Corninrmus Filament
Reinforced W/NiAl and A.120@iAl timposiles”, Submittrd for publication.