Interface Composite

7172019 Interface Composite

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In1047298uence of interfaces on the mechanical behavior of SiCparticulate-reinforced AlndashZnndashMgndashCu composites

Jingya Song a Qiang Guo an Qiubao Ouyang a Yishi Su a Jie Zhang a Enrique J Lavernia b Julie M Schoenungb Di Zhang an

a State Key Laboratory of Metal Matrix Composites Shanghai Jiao Tong University 800 Dongchuan Road Shanghai 200240 Chinab Department of Chemical Engineering and Materials Science University of California Davis CA 95616 USA

a r t i c l e i n f o

Article history

Received 29 April 2015

Received in revised form

17 July 2015

Accepted 17 July 2015Available online 18 July 2015

Keywords

Metal matrix composites

Nanoindentation

Aging

Interface structure

Dislocation distribution

a b s t r a c t

In particulate-reinforced metal matrix composites (MMCs) geometrically necessary dislocations (GNDs)

form in the vicinity of reinforcementmatrix interfaces In this study the hardness distribution across the

interface was studied using nanoindentation with high spatial resolution for composites treated under

different aging conditions The size of the GND punched zone as determined from the hardness mea-

surement was found to be in agreement with that estimated by transmission electron microscopy (TEM)

Mechanical characterization of bulk composites revealed a reduction in failure strain with decreasing

punched zone size while the strength of the composites was found to depend more on the intrinsic

strength of the matrix alloy These observations were interpreted in terms of the load transfer capacity

between the matrix and reinforcement through the interface

amp 2015 Elsevier BV All rights reserved

1 Introduction

Particulate-reinforced metal matrix composites (MMCs) are

ideally suited for many structural and functional applications be-

cause of their high speci1047297c strength and stiffness isotropic prop-

erties and relatively simple processing as compared with mono-

lithic materials and conventional 1047297ber-reinforced composites [1ndash

4] Among the parameters that may affect the mechanical perfor-

mance of these composites such as the reinforcement particle size

[56] distribution [7] and volume fraction [38] the properties at

and in the vicinity of particlendashmatrix interfaces play a critical role

[910] A strong interface would usually allow for effective load

transfer from the matrix to the reinforcement leading to improved

strength stiffness and resistance to environmental attack [11] The

interfacial properties have also been found to determine the fail-ure mode of the composite [124] where failure initiated by in-

terfacial debonding is likely to occur when the interface is weak

During the processing of particulate-reinforced MMCs geo-

metrically necessary dislocations (GNDs) typically form in the

metal matrix close to the particlendashmatrix interface as a result of

the residual stress caused by the mismatch in the coef 1047297cients of

thermal expansion (CTE) between particle and matrix [25] The

presence of GNDs which reportedly hardens the matrix alloy inthe vicinity of the interface has been con1047297rmed experimentally

[16] as well as predicted on the basis of numerical models [512]

The nature of the regions populated with GNDs around the par-

ticle in the metal matrix (dislocation punched zone) is depen-

dent on the size morphology and distribution of the reinforce-

ment particles [513] as well as on the processing and heat

treatment parameters [14] The in1047298uence of these regions on

mechanical behavior is twofold On one hand these regions are

considered to provide the main strengthening mechanism in

particulate-reinforced MMCs [51415] On the other hand they

also give rise to rapid hardening saturation of the matrix at low

external strain values resulting in the degradation in the compo-

sites ductility and fracture toughness [1016] In addition to their

in1047298uence on mechanical properties for composites whose matrixalloy is age-hardenable the existence of the GNDs and the asso-

ciated dislocation-punched zones have been reported to sig-

ni1047297cantly facilitate the aging kinetics of the composites during

heat treatment where the GNDs were proposed to serve as het-

erogeneous nucleation sites for precipitate formation [17]

Although extensive experimental and theoretical research have

been carried out to study the size and distribution of dislocation

punched zones as a function of the con1047297guration of the re-

inforcement particles and their effect on the aging response of the

composites inspection of the published literature show that there

are limited studies devoted to clarifying how such regions would

Contents lists available at ScienceDirect

journal homepage wwwelseviercomlocatemsea

Materials Science amp Engineering A

httpdxdoiorg101016jmsea201507050

0921-5093amp 2015 Elsevier BV All rights reserved

n Corresponding authors

E-mail addresses guoqsjtueducn (Q Guo) zhangdisjtueducn (D Zhang)

Materials Science amp Engineering A 644 (2015) 79ndash84





negligible effect on indentation hardness In general a shallower

indentation would inevitably result in a smaller lateral dimension

of the indents allowing for a smaller indent spacing and subse-

quently a more accurate estimate on the hardness variation and

punched zone size However owing to the roughness of the po-lished sample surface it was found that for indentation depths

smaller than 100 nm good convergence on the calculated hardness

was not achieved and accordingly a 1 mm spacing between

neighboring indents and a 120 nm maximum indentation depth

were used in the following experiments to ensure a high spatial

resolution of the indents as well as the representativeness of the

trend in hardness across the interface as illustrated in Fig 2 For

each sample at least 4 arrays of indentation across different SiC

particles were conducted and more than 20 individual indents

were made for each indentation array covering the entire range

from the SiC particle (characterized by a plateau in hardness cor-

responding to SiC) across the SiCAl interface and 1047297nally to the

matrix alloy (characterized by a plateau in hardness corresponding

to the matrix)Fig 3(a) shows a set of representative data corresponding to

the hardness distribution across the interface for samples treated

under 4 different processing conditions (as-extruded aged at 12

24 and 48 h) The origin on the horizontal axis ( xfrac140) corresponds

to the approximate position of the interface right of which ( x40)

corresponds to the Al matrix and left of which ( xo0) corresponds

to the SiC particle In the xo0 regime (particle) the hardness

plateau has a magnitude of 40 GPa corresponding to that of the

particle In the x40 regime (matrix) the magni1047297ed view as shown

in Fig 3(b) demonstrates that for the indentations made more

than 6 mm away from the SiCAl interface the hardness values

for all the samples reach a plateau corresponding to the intrinsic

hardness of the matrix alloy In this region the sample aged for

24 h possesses the highest hardness of 2867016 GPa followed

by the sample aged for 48 h (2727015 GPa) the sample aged for

12 h (2237001 GPa) and the as-extruded sample

(1967004 GPa) Therefore the samples aged for 12 24 and 48 h

are hereafter denoted as underaged peak-aged and overagedsamples respectively Fig 3(b) also demonstrates that for the as-

extruded and underaged samples the hardness generally shows a

gradual reduction from the interface towards the matrix and the

width of this transition region can be estimated to be 6 mm for

both cases For the peak-aged and overaged counterparts in

comparison the transition is rather sharp and the widths are

shown to be smaller than 2 mm ie comparable to the spacing

between neighboring indents The punched zone size together

with the hardness of the matrix under different aging treatments

is documented in Table 1

32 Microstructure evolution during aging process

It has been proposed that the width of the hardness transitionregion is a measure of the dislocation punched zone size asso-

ciated with the GNDs formed upon the processing of metal matrix

composites [15] Therefore a high density of dislocations is ex-

pected to be present in this region which is exactly what has been

observed Fig 4 shows the representative TEM images on the SiC

Al interface taken under two beam conditions from [011] zone

axis and [002] diffraction vector for the 4 sets of samples pro-

cessed under different heat treatments At least 3 TEM images

were used for each sample to estimate the width of its corre-

sponding punched zone size As demonstrated the dislocation

punched zone extends 2ndash4 mm into the matrix alloy for composites

in the as-extruded and underaged states and are limited to 1ndash

2 mm for the peak-aged and overaged samples This is in qualita-

tive agreement with the results obtained from the hardnessmeasurement However it should be noted that the punched zone

size estimated from the TEM images only represents its lower limit

because not all dislocations are visible under this particular dif-

fraction condition This is especially true for the as-extruded

sample where the size of the grains immediately adjacent to the

SiC particles was only about 2 mm (Fig 4a) making the dislocations

in the neighboring grains unobservable so its punched zone size

may be substantially underestimated Also careful examination at

the SiCAl interfaces in Fig 4 reveals that dislocations adjacent to

Fig 2 SEM image of an indentation array made in the 24 h-aged composite with

120 nm indentation depth and 1 μm indentation spacing

Fig 3 (a) Variation of indentation hardness across the SiCpAl interface for composites treated under 4 different processing conditions (b) zoomed-in rendition of the boxed

region in Fig (a)

J Song et al Materials Science amp Engineering A 644 (2015) 79ndash84 81



the interface are not uniformly distributed and regions of slightly

different diffraction contrast can be observed In general micro-

structural features such as grain boundaries interfaces pre-

cipitates and stacking faults can all affect the distribution of dis-

locations and subsequently the formation and morphology of

substructures On the other hand as mentioned previously dis-

locations have been reported to serve as heterogeneous nucleation

sites for precipitate formation in age-hardenable alloys [17]

Therefore the presence of SiCAl interfaces as well as the pre-

cipitates introduced by aging may have given rise to substructure

formation (ie the different diffraction contrast) in the vicinity of

the interface as observed in Fig 4 Moreover it is important topoint out that the indentation depth (120 nm) and the associated

lateral radius of the indents (500 nm [23]) is signi1047297cantly (at

least a factor of 8) smaller than the average grain size of the Al

matrix in all samples (4ndash65 μm see Table 1 Column 8) Yang and

Vehoff [24] reported that if the grain size of pure Ni is more than

35 times the lateral radius of the indents the plastic zone caused

by the indentation conducted at the center of the grain would be

limited inside the grain and would not interact with the grain

boundaries This is generally the case in the present experiments

Therefore on the whole the hardness evolution in the vicinity of

the SiCAl interface was unlikely to be caused by the variation in

the Al grain size and should solely be the result of the GNDs

During cooling from the high temperature processing of stir

casting and extrusion deformation the CTE mismatch between the

particle and the matrix alloy is likely to cause residual stress in the

vicinity of the interface [1] and subsequently a considerable in-

crease in the dislocation punched zone size in the as-extruded and

underaged samples On the other hand the reduction in the

punched zone size along with the progression of precipitation in

the peak-aged and overaged samples can be explained in terms of

the aging kinetics affected by the presence of GNDs It is wellknown that in an AlndashZnndashMgndashCu alloy the precipitation sequence

upon arti1047297cial aging can be generally listed as GuinierndashPreston

(GP) zones-ηrsquo and η (MgZn2)-S (Al2CuMg) phases [25] In re-

lated studies Cantrell et al [26] argues that the nucleation and

growth of precipitates near matrix dislocations may cause the

dislocations to bow under the action of local precipitatendashmatrix

coherency stresses it has also been reported that when the pre-

cipitates become larger in size with increasing aging time they

may promote the rearrangement of dislocations and the

Table 1

Tensile properties indentation properties and structural parameters of the SiCpAl composite samples treated under different processing conditions

Condition Punched zone size (μm) Matrix hardness (GPa) Strength (MPa)a ΔsL-T (MPa) Failure strain () Average grain size (μm) ΔsL-T s y ()

s y smatrix

As-extruded 6 1967004 370 345 25 62 4 68

Underaged 6 2237001 480 393 87 43 51 181

Peak-aged o2 2867016 520 503 17 26 27 33

Overaged o2 2727015 496 478 18 19 65 36

as y 02 offset yield strength of the composite smatrix 02 offset yield strength of the Al matrix

Fig 4 Transmission electron microscopy (TEM) images of the reinforcementmatrix interface from SiCpAl composite samples treated under different processing condi-

tions (a) as-extruded state (b) underaged state (c) peak-aged state and (d) overaged state

J Song et al Materials Science amp Engineering A 644 (2015) 79ndash8482



movement of subgrain boundaries [27] Furthermore GNDs have

been proposed to signi1047297cantly facilitate the aging kinetics by ser-

ving as heterogeneous nucleation sites for precipitate formation

[2] All these effects may lead to a decrease in the GND density and

the associated punched zone size as has been observed here in the

TEM images

EBSD measurements show that the average grain size of the

matrix alloy changes from 51 μm for the underaged sample down

to 27 μm for the peak-aged sample and again up to 65 μm for theoveraged sample (Table 1) This evolution was mostly likely the

result of recrystallization upon aging treatment Recrystallization

of deformed alloys at aging temperatures has been reported but

the kinetics can be quite slow For example Jones and Hymphreys

[28] reported the complex interaction between precipitate for-

mation and recrystallization in AlndashSc alloys where precipitation

was found to precede follow or occur concurrently with re-

crystallization depending on the deformation processing condi-

tions of the alloy It was also found that it takes over 20 h for the

Alndash002 Sc wt alloy to get fully recrystallized at 280 degC Liu et al

[27] reported recrystallization during aging for a rapidly solidi1047297ed

CundashCrndashZrndashMg alloy and the dispersed precipitates were reported

to retard recrystallization In the solution heat-treated Mgndash9Alndash

1Zn alloy which was later subject to severe plastic deformationrecrystallization was observed with suf 1047297cient aging time (410 h)

at 180 degC [29]

Considering the evolution of the average grain size of the Al

matrix as demonstrated in Table 1 we propose that the nucleation

and growth of recrystallized grains started to take place at some

point between the underaging and peak-aging processes ie be-

tween 12 h and 24 h aging treatment During aging towards the

overaged state further growth of recrystallized grains occurred

leading to an even larger average grain size of 65 μm On the other

hand the observation that the average grain size changed from

4 μm in the as-extruded state to 51 μm in the underaged state is

likely the result of signi1047297cant grain growth caused by solution

heat-treatment [3031]

Since recrystallization (which softens the matrix) was likely to

occur concurrently with aging-induced precipitation (which

hardens the matrix) in this study the resulting hardness evolution

in the Al matrix in the vicinity of the interface would be de-

termined by the competition between these two processes Results

of hardness measurements as well as macroscopic tensile tests

indicate that precipitation-strengthening played the dominant

role making the 24 h-aged (ldquopeak-agedrdquo) sample stronger and

harder than the 12 h-aged (ldquounderagedrdquo) sample

33 Mechanical properties of the bulk composites

Fig 5 shows the engineering stress vs engineering strain re-

sponse of the composites as obtained from uniaxial tensile tests

Consistent with the hardness measurements the sample subjected

to the 24 h aging treatment (peak-aged) has the highest 02yield strength (520 MPa) followed by the overaged (496 MPa)

underaged (480 MPa) and as-extruded (370 MPa) samples In the

case of ductility the failure strain shows a decrease with in-

creasing aging time changing from 62 for the as-extruded

composite to 43 (underaged) 26 (peak-aged) and 1047297nally to

19 (overaged) The yield strength and failure strain data are lis-

ted in Table 1 along with the estimated dislocation punched zone

size and the intrinsic hardness of the matrix alloy obtained from

hardness measurements and the average grain size of the matrix

alloy for each sample measured by EBSD characterization

34 Strengthening mechanisms of the composites

To explore the mechanisms that may be responsible for the

evolution of the compositesrsquo strength and failure strain for differ-

ent heat treatments the yield strength of the composites s y isdecomposed and expressed as [710]

1

y aging SiC

aging dis L T

matrix L T

0

0

σ σ σ σ

σ σ σ σ

σ σ

= + Δ + Δ

asymp + Δ + Δ + Δ

asymp + Δ ( )

minus

minus

where s0 stands for the initial strength of 7A04 alloy without re-

inforcement and aging treatment Δsaging and ΔsSiC are the

strengthening contributions from arti1047297cial aging and SiC re-

inforcement respectively ΔsSiC primarily comes from a disloca-

tion strengthening part Δsdis and a load-transfer (L-T ) strength-

ening part ΔsL-T where the former is usually negligible for mi-

cron-sized reinforcements [7] matrix aging 0σ σ σ = + Δ is the strength

of the matrix alloy which can be estimated from the hardnessplateau corresponding to the matrix alloy by nanoindentation

measurement (Fig 3 and Column 3 in Table 1)

A Tabor factor value of 3 is usually used to estimate the yield

strength from hardness data [32] ie 3 yσ Η asymp where H is the

hardness measured by nanoindentation However in the present

work dividing the hardness of 286 GPa for the peak-aged com-

posite (Table 1) by 3 gives a stress value of about 950 MPa which

is almost a factor of 2 higher than the yield strength of monolithic

7A04 Al alloy fabricated and heat treated under nominally the

same conditions as those used in this study (503 MPa) [33] A si-

milar discrepancy has previously been reported for 6xxx Al alloys

[1334] and is likely the result of strong dislocation pile-up at the

indents Bolshakov and Pharr [35] reported that for materials

showing a strong pile-up effect (a high h f hmax where h f is the 1047297nalcontact depth and hmax is the maximum contact depth during the

nanoindentation test [2135]) because of a signi1047297cant under-

estimation of the indentation contact area the hardness calculated

using the OliverndashPharr method [21] can be greatly overestimated

When h f hmax407 the underestimate of contact area can reach as

high as 60 [35] In this study h f hmax is found to be systematically

higher than 09 indicating a strong pile-up and subsequently a

substantial overestimation of the hardness values To accom-

modate this effect here we de1047297ne an ldquoeffective Tabor factorrdquo

calculated using the ldquoapparentrdquo hardness of 7A04 alloy matrix

(286 GPa see Table 1) divided by the yield strength of peak-aged

7A04 alloy (503 MPa) [33] which gives a value of 568 The yield

strength of the matrix alloy in the composites smatrix can then be

estimated from the hardness data and the load-transfer

Fig 5 Tensile stress vs strain curves of the composite samples treated under

4 different processing conditions




strengthening contribution from the SiC particles ΔsL-T can be

calculated from Eq (1) as matrix0σ σ minus Column 9 in Table 1 lists the percentage of ΔsL-T in s y for the

samples heat treated under different conditions It is clearly shown

that samples with larger dislocation punched zones (the as-ex-

truded and underaged states) have relatively more effective load

transfer between the reinforcement and the matrix while for the

ones with narrower punched zones (the peak-aged and overaged

states) the load carried by the reinforcement is very limitedwhich would more easily cause stress concentration and sub-

sequent necking upon deformation leading to early failure of the

composite [20] This argument is supported by our results which

show a decreasing failure strain from the as-extruded and un-

deraged samples relative to the peakaged and overaged ones The

reduction in ductility may also be associated with different types

and contents of the precipitates in the samples whose interface

with the matrix alloy may serve as nucleation sites and easy paths

for crack growth [36] In the case of the compositesrsquo strength on

the other hand it seems that the intrinsic strength of the matrix

alloy plays a dominant role over the interfacial properties and the

composite with the strongest matrix (peak-aged) has the highest

yield strength This is actually a natural result of the low load

transfer in the interface region for this sample

4 Conclusions

This work studied the hardness distribution across the re-

inforcementmatrix interface in SiCpAl composites treated at

different aging conditions with high spatial resolution It has been

found that the dislocation punched zone size determined from the

hardness measurements is in qualitative agreement with the range

that GNDs extended out from the interface measured in TEM mi-

crostructural analysis The evolution of the width of the punched

zone over different heat treatment conditions can be rationalized

by the effect of GNDs on aging kinetics Moreover mechanical

characterization of bulk composites revealed a reduction in failure

strain with decreasing punched zone size most likely due to theeffect of less effective load transfer between the reinforcement and

the matrix in composites with smaller punched zone sizes In the

case of composite strength it has been found that the strength is

more dependent on the intrinsic strength of the matrix alloy then

on the nature of the interface This study indicates that the

hardness measurement combined with microstructural analysis

may have important implications on the bulk properties of parti-

culate-reinforced MMCs and thus be a useful way to help sort out

composites with speci1047297c properties leading to improved modeling

and design of MMCs

Acknowledgment

The work is supported by the National Basic Research Program

of China (973 Program No 2012CB619600) the National High-

Tech RampD Program of China (863 Program No 2013AA031201) the

National Natural Science Foundation of China (Program No

51471190) the Science and Technology Commission of Shanghai

Municipality (STCSM) (Nos 13PJ1404000 and 14DZ2261200) and

the Of 1047297ce of Naval Research (ONR) under the guidance of Dr

Lawrence Kabacoff (ONR N00014-12-1-0237) QG and JS would

like to acknowledge C Xie from Keysight Technologies (China)

and Dr L Jiang from the University of California Davis for useful

discussions

References

[1] N Chawla YL Shen Adv Eng Mater 3 (2001) 357ndash370[2] S Suresh KK Chawla Fundamentals of Metal Matrix Composites Butter-

worth-Heinemann Stoneham MA 1993[3] T Ozben E Kilickap O Ccedilakır J Mater Process Technol 198 (2008) 220ndash225[4] SG Song N Shi GT Gray III JA Roberts Metall Mater Trans A 27 (1996)

3739ndash3746[5] YS Suh SP Joshi KT Ramesh Acta Mater 57 (2009) 5848ndash5861[6] N Chawla C Andres JW Jones JE Allison Metall Mater Trans A 29 (1998)

2843ndash2854[7] X Kai Z Li G Fan Q Guo D-B Xiong W Zhang Y Su W Lu WJ Moon

D Zhang Mater Sci Eng A 587 (2013) 46ndash53[8] X Zhang L Geng G Wang J Mater Process Technol 176 (2006) 146ndash151

[9] Z Wang M Song C Sun D Xiao Y He Mater Sci Eng A 527 (2010)6537ndash6542

[10] R Vogt Z Zhang Y Li M Bonds ND Browning EJ Lavernia JM SchoenungScr Mater 61 (2009) 1052ndash1055

[11] JK Park JP Lucas Scr Mater 37 (1997) 511ndash516[12] J Shao B Xiao Q Wang Z Ma K Yang Compos Sci Technol 71 (2011) 39ndash45[13] C Liu S Qin G Zhang M Naka Mater Sci Eng A 332 (2002) 203ndash209[14] J Ye J He JM Schoenung Metall Mater Trans A 37 (2006) 3099ndash3109[15] S Qin C Liu J Chen G Zhang W Wang J Mater Sci Lett 18 (1999)

1099ndash1100[16] X Kai Z Li G Fan Q Guo Z Tan W Zhang Y Su W Lu WJ Moon D Zhang

Scr Mater 68 (2013) 555ndash558[17] A Deschamps Y Brechet Acta Mater 47 (1999) 293ndash305[18] Q Ouyang R Li W Wang G Zhang D Zhang Mater Sci Forum 546ndash549

(2007) 1551ndash1554[19] D Ge V Domnich T Juliano EA Stach Y Gogotsi Acta Mater 52 (2004)

3921ndash3927[20] G Meijer F Ellyin Z Xia Compos Part B ndash Eng 31 (1999) 29ndash37[21] WC Oliver GM Pharr J Mater Res 7 (1992) 1564ndash1583[22] A Poudens B Bacroix T Bretheau Mater Sci Eng A 196 (1995) 219ndash228[23] M Goumlken M Kempf Acta Mater 47 (1999) 1043ndash1052[24] B Yang H Vehoff Acta Mater 55 (2007) 849ndash856[25] MM Sharma MF Amateau TJ Eden Mater Sci Eng A 424 (2006) 87ndash96[26] JH Cantrell WT Yost Appl Phys Lett 77 (2000) 1952ndash1954[27] P Liu BX Kang XG Cao JL Huang B Yen HC Gu Mater Sci Eng A 265

(1999) 262ndash267[28] MJ Jones FJ Humphreys Acta Mater 51 (2003) 2149ndash2159[29] YC Yuan AB Ma JH Jiang Y Sun FM Lu LY Zhang D Song J Alloy

Compd 594 (2014) 182ndash188[30] YH Zhao XZ Liao S Cheng E Ma YT Zhu Adv Mater 18 (2006)

2280ndash2283[31] H Li QZ Mao ZX Wang FF Miao BJ Fang ZQ Zheng Mater Sci Eng A 620

(2015) 204ndash212[32] JR Cahoon WH Broughton AR Kutzak Metall Trans 2 (1971) 1979ndash1983[33] Y Su Q Ouyang W Zhang Z Li Q Guo G Fan D Zhang Mater Sci Eng A

597 (2014) 359ndash369[34] S Qin C Chen G Zhang W Wang Z Wang Mater Sci Eng A 272 (1999)

363ndash370[35] A Bolshakov GM Pharr J Mater Res 13 (1998) 1049ndash1058[36] YH Zhao XZ Liao Z Jin RZ Valiev YT Zhu Acta Mater 52 (2004)

4589ndash4599

J Song e t al Materials Science amp Engineering A 644 (2015) 79ndash8484











































































region in Fig (a)












































Table 1



s y smatrix

As-extruded 6 1967004 370 345 25 62 4 68

Underaged 6 2237001 480 393 87 43 51 181

Peak-aged o2 2867016 520 503 17 26 27 33

Overaged o2 2727015 496 478 18 19 65 36












TEM images

















at 180 degC [29]







































1

y aging SiC

aging dis L T

matrix L T

0

0

σ σ σ σ

σ σ σ σ

σ σ

= + Δ + Δ


asymp + Δ ( )

minus

minus





























































4 Conclusions




















and design of MMCs

Acknowledgment











discussions

References



















4589ndash4599










































































region in Fig (a)












































Table 1



s y smatrix

As-extruded 6 1967004 370 345 25 62 4 68

Underaged 6 2237001 480 393 87 43 51 181

Peak-aged o2 2867016 520 503 17 26 27 33

Overaged o2 2727015 496 478 18 19 65 36












TEM images

















at 180 degC [29]







































1

y aging SiC

aging dis L T

matrix L T

0

0

σ σ σ σ

σ σ σ σ

σ σ

= + Δ + Δ


asymp + Δ ( )

minus

minus





























































4 Conclusions




















and design of MMCs

Acknowledgment











discussions

References



















4589ndash4599












































Table 1



s y smatrix

As-extruded 6 1967004 370 345 25 62 4 68

Underaged 6 2237001 480 393 87 43 51 181

Peak-aged o2 2867016 520 503 17 26 27 33

Overaged o2 2727015 496 478 18 19 65 36












TEM images

















at 180 degC [29]







































1

y aging SiC

aging dis L T

matrix L T

0

0

σ σ σ σ

σ σ σ σ

σ σ

= + Δ + Δ


asymp + Δ ( )

minus

minus





























































4 Conclusions




















and design of MMCs

Acknowledgment











discussions

References



















4589ndash4599









TEM images

















at 180 degC [29]







































1

y aging SiC

aging dis L T

matrix L T

0

0

σ σ σ σ

σ σ σ σ

σ σ

= + Δ + Δ


asymp + Δ ( )

minus

minus





























































4 Conclusions




















and design of MMCs

Acknowledgment











discussions

References



















4589ndash4599

























4 Conclusions




















and design of MMCs

Acknowledgment











discussions

References



















4589ndash4599


Interface Composite

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