agrawal, p. - fracture in metal–ceramic composites

8/2/2019 Agrawal, P. - Fracture in metalceramic composites

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Fracture in metalceramic composites

Parul Agrawal *,1, C.T. Sun

School of Aeronautics and Astronautics, Purdue University, West Lafayette, IN 47907-1282, USA

Received 20 August 2003; received in revised form 21 August 2003; accepted 22 September 2003

Available online 15 December 2003

Abstract

This research focuses on fracture mechanisms in metalceramic composites. Two co-continuous composites, Cu/Al 2O3 and Al/

Al2O3 and a metalmatrix composite Al/SiC were studied. It was found that each composite displayed a different fracture mech-

anism. The crack propagation inside the metalmatrix composite was dominated by the Al matrix characteristics. However, crack

propagation inside both the co-continuous composites was influenced by their microstructure, thermal residual stresses and con-

tiguity. This unique fracture characteristic of co-continuous composites has been elucidated in the present research by experi-

mentation as well as computational modeling. In situ three-point bend tests were performed inside an environmental scanning

electron microscope chamber to observe crack growth at the microstructural scale. Finite element modeling was performed by using

globallocal approach to simulate crack propagation and understand the effects of the microstructure and thermal residual stresses.

It was shown that the crack propagated inside the metallic phase and at the interface for the Cu/Al 2O3 composite due to a high level

of tensile thermal stresses inside the metallic phase, as well as due to low contiguity of ceramic phase. However, in the case of Al/

Al2O3 composite, the crack propagated inside the ceramic due to significantly smaller thermal stresses inside the metallic phase as

well as higher contiguity of ceramic phase.

2003 Elsevier Ltd. All rights reserved.

Keywords: B. Fracture; B. Interface; C. Crack; D. Scanning electron microscopy; Co-continuous composites

1. Introduction

Metalceramic composites have been a topic of

interest for many researchers for various reasons [15].

The motivation for developing metalceramic compos-

ites is to fabricate structures that possess superior stiff-

ness compared to metals, simultaneously having better

toughness and structural integrity compared to a

monolithic ceramic. The metalceramic composites that

consist of interconnecting network of metal and ceramic

phases are defined as co-continuous metalceramic

composites. The current study is focused on the fracture

mechanisms inside these composites. A metalmatrix

composite with ceramic reinforcements, Al/SiC, has also

been investigated to compare and contrast the fracture

mechanisms inside the two categories of metalceramic

composites.

Sufficient literature can be found for the fracture of

metalmatrix composites with reinforced ceramic parti-

cles or fibers. Only the metal phase is contiguous in these

composites. As a result, yielding and fracture is domi-

nated by the metal phase. Foo and coworkers [1] focused

on interface characterization and the effects of interfacial

strength on failure and debonding in particulate and

whisker reinforcedcomposites. Davidson and Regener [2]

performed in situ tensile tests inside a scanning electron

microscope (SEM) chamber and recorded the failure

mechanisms inside coated and uncoated Al/SiC particu-

latecomposites. Suery and LEsperance [3] andMcDanels

[4] also studied Al/SiC composites.

At the other end of the spectrum, some studies were

performed on ceramicmatrix composites with partic-

ulate metallic phase. Inclusion of ductile particles in a

brittle matrix leads to toughening. The main toughen-

ing mechanism reported was crack bridging. This

model was first introduced by Krstic [5]. It was

Composites Science and Technology 64 (2004) 11671178

www.elsevier.com/locate/compscitech

COMPOSITES

SCIENCE AND

TECHNOLOGY

* Corresponding author. Tel.: +1-408-256-5832; fax: +1-408-256-

2410.

E-mail address: [email protected] (P. Agrawal).1 Present address: Hitachi (Previously IBM), 5600 Cottle Road, San

Jose, CA 95193, USA.

0266-3538/$ - see front matter 2003 Elsevier Ltd. All rights reserved.

doi:10.1016/j.compscitech.2003.09.026
http://mail%20to:%[email protected]/http://mail%20to:%[email protected]/


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reported by Toya [6] that the fracture toughness of

metal reinforced composites was proportional to the

ductile particle size until a critical point. Beyond that

critical size thermal residual stresses were found to play

a role in weakening the interface. Lange [7] proposed a

fracture mechanics model to explain the effect of par-

ticle size on interface debonding due to thermal stres-ses. Kohle et al. [8] performed modeling to predict the

critical size of Ni particles in the Ni/Al2O3 composite

beyond which crack extension will start at the interface

due to thermal residual stresses. The crack tip bridging

mechanism by ductile particles is also discussed by Ohji

et al. [9] for ceramic nanocomposites. However, the

fracture behavior and failure mechanisms inside co-

continuous ceramic material has not been investigated

by many researchers to authors knowledge. Most of

the researchers have investigated either metalmatrix or

ceramicmatrix composites. Cichocki [10] performed

bend tests to understand the failure mechanisms inside

co-continuous composites and relate it to their micro-

structure. In the present paper, the fracture in two

types of co-continuous composites Cu/Al2O3 and Al/

Al2O3 is described and compared with a metalmatrix

composite Al/SiC.

To observe crack propagation and fracture charac-

teristics three-point bend tests were performed inside an

environmental scanning electron microscope (ESEM)

chamber. These experiments enabled an observation and

understanding of the various failure mechanisms that

take place during crack propagation. It also provided an

opportunity to understand the role of various factors

like residual stresses, grain boundaries and contiguity,inside a multiphase system.

In order to understand and explain the failure

mechanisms and experimental results, finite element

simulations were performed by using the software

package, Franc 2D, developed by Wawrzynek and

Ingraffea from Cornell University [11]. This software

enables adaptive remeshing during crack propagation.

Hence, it enabled the authors to determine the stress

distribution, stress intensity factors and failure mecha-

nisms that were observed in three-point bend experi-

ments in terms of residual stresses.

2. Processing and microstructure of the composites

2.1. Processing

The two co-continuous composites were made by

metal infiltration techniques inside Al2O3 ceramic

sponges. In the case of Cu/Al2O3 composite, molten

copper alloy was infiltrated inside a sponge made of

spherical Al2O3 particle of 5 lm radius that created a co-

continuous network of metal and ceramic phase. Some

of the processing details are provided by Gonzales and

Trumble [12]. We were able to create a uniform co-

continuous network of Cu and Al2O3.

Al/Al2O3 composite was prepared by a fugitive sin-

tering process to provide spherical aluminum metal

network, co-continuous with Al2O3 network. The Al2O3grains were very small (0.3 lm) for this composite. The

processing details were provided by Cichocki [10]. Themetalmatrix composite was provided by ALYN Cor-

poration, Irvine, CA. This composite was processed by a

powder metallurgy process. 6092 Aluminum alloy was

reinforced with SiC particles.

2.2. Microstructure

Fig. 1 shows the SEM micrograph of the Cu/Al2O3composite. There is a continuous network of metal in-

terconnected with a continuous network of the ceramic

phase. This micrograph is a backscattered image. The

dark phase is the ceramic and the brighter phase cor-

responds to copper. XRD spectra confirmed that only

Al2O3 and Cu alloy phases exist in the composite and no

significant reaction had taken place. XRD data also

showed that no anisotropy is present in the Al 2O3 pre-

forms. The continuity of the metal phase was confirmed

by the electrical conductance tests. Further details are

provided by Agrawal [13].

Fig. 2 shows an optical micrograph of Al/Al2O3composite. The lighter phase is metallic Al and the

darker phase is Al2O3. The Al particles are about 100

200 lm in diameter. The grain size of Al2O3 particles

used was 0.31.0 lm. The contiguity of the Al2O3 phase

was found to be greater in this composite compared tothat in the Cu/Al2O3 composite. The volume fraction of

metal Al was computed as 70% and that of Al2O3 is

30%. The continuity of metal phase was again checked

by the electrical conductance tests.

Fig. 1. SEM image of Cu/Al2O3 composite at 2300 magnification.

1168 P. Agrawal, C.T. Sun / Composites Science and Technology 64 (2004) 11671178


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An optical micrograph of 6092 Al/SiC composite isshown in Fig. 3. The dark phase corresponds to SiC

ceramic and the lighter phase corresponds to the Al

matrix. It is evident from the micrograph that the metal

phase is continuous and that the ceramic particles do

not form a contiguous network, i.e., the particles are

suspended in the matrix phase.

Physical properties like volume fraction, contiguity of

the spherical phase (Al2O3 in case of Cu/Al2O3 com-

posite and Al in case of Al/Al2O3), Youngs moduli of

the co-continuous composites are provided in Table 1.

The details of the measurements and computations of

physical properties are provided by Agrawal [13].

3. Thermal residual stress

Thermal residual stresses are generated in metal

ceramic composites during cooling after significantly

high (12001300 C) processing temperature. The mis-

match in the coefficient of thermal expansion between

metal and ceramic phases causes these stresses. After the

cooling process, the metal and ceramic phases of the

composite develop tensile and compressive stresses,

respectively. These stresses affect failure mechanisms

inside the composites. One of the objectives behind the

present research was to investigate the effect of thermal

stresses on failure mechanisms of metalceramic com-

posites. The non-destructive technique of neutron dif-

fraction technique was adopted to measure thermal

stresses in Cu/Al2O3 and Al/Al2O3 composites. The ex-

periments were performed at Chalk River National

Laboratory in Ontario, Canada. The composites were

bombarded with monochromatic thermal neutron

beams, which penetrated through the thickness of thespecimen. By measuring the shifts in spectral peaks of

metal and ceramic phases, the strains and hence thermal

residual stresses, in each phase were obtained. Details of

these experiments and analysis are provided by Agrawal

et al. [14]. For the Cu phase in the Cu/Al 2O3 composite,

the stress measurements were done with respect to peak

shifts in [2 0 0] and [3 1 1] planes. The computed stress

tensors were as follows:Fig. 3. Optical micrograph of 6092 T6 Al/SiC composite at 1000magnification.

Table 1

Material properties used in the simulations

Material Volume

fraction

Contiuity

of spherical

phase

Density

(kg/m3)

E (GPa) Poissons

ratio, t

Fracture

toughness

KIC(MPa

ffiffiffiffim

p)

Coefficient

of thermal

expansion

106/C (a)

Effective process

temperature DT

(C)

Cu/Al2O3 30% Cu, 70%

Al2O3

0.5 5340 250.0 0.31 3.8 8.95 700

Al/Al2O3 70% Al, 30%

Al2O3

0.3 3080 120.0 0.3 6.5 16.8 150

Cu alloy 8960 114.0 0.33 12.0 17.3

Al metal 2700 70.0 0.35 29.0 23.7

Al2O3 3900 380.0 0.22 4.5 6.6

Fig. 2. Optical micrograph of Al/Al2O3 composite at 100magnification.

P. Agrawal, C.T. Sun / Composites Science and Technology 64 (2004) 11671178 1169


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rm2 0 0 578 17 817 456 218 21 526

264

375 MPa;

rm3 1 1 720 8 128 712 1212 12 712

2

64

3

75 MPa:

For the Al2O3 phase in Cu/Al2O3 composites, the fol-

lowing values were computed corresponding to [2 1 9]

and [3 0 0] planes:

rc2 1 9 155 3 6

3 164 46 4 179

264

375 MPa;

rc3 0 0 217 4 04 202 1

0 1

226

2

64

3

75MPa:

For the Al/Al2O3 composite, the tensile stress tensor for

the metal phase was

rm2 2 0 110 1 3

1 113 4

3 4 113

24

35 MPa:

For the Al2O3 phase in Al/Al2O3 composite,

rc2 1 3 205 6 6

6 156 16 1

159

264

375

MPa;

rc2 1 6 164 18 21

18 149 1821 18 185

264

375 MPa:

The diffraction experiment confirmed that, for both

composites, the stresses in the metal and ceramic phases

are tensile and compressive in nature, respectively. It

also shows that the hydrostatic component of the stress

tensors dominate in both cases. This is because the

technique provides an average measure of stresses over

the entire volume of the specimen. In the case of Cu/

Al2O3 composite the stress level in the metal phase was

higher (about 620 MPa) as compared to that in the Al/

Al2O3 composite (100 MPa), even though the com-

pressive stresses in the Al2O3 phases turned out to be

very close to each other. The volume fraction difference

in the two composites and the lower melting point of Al

metal causes this difference. These stresses influence

fracture properties of the composites to a great extent,as shown in the later sections.

4. Specimen preparation and testing procedure for ESEM

experiments

ASTM standards E399 [15] were used to provide a

general guideline in specimen preparation. The com-

posites were cut to the desired length and width with a

diamond wheel on a surface grinder. Single edge V

notched beam (SEVNB) specimens were prepared for

the three-point bend fracture tests. Table 2 lists the

specimen dimensions used for fracture tests. The speci-

men surfaces were polished to a very fine finish. The

polishing was performed in several stages, starting with

380 coarse grit sand paper, then going to 400, 500, 600

and 1200 grits. A final finish was given with nylon cloth

in 0.05 lm colloidal silica medium.

Al/Al2O3 samples were etched for 1 min in sodium

hydroxide. This process removed about 2550 lm thin

Al metal layer from the surface to give enough surface

geometry for secondary electron emission. In the case of

Cu/Al2O3 composites, backscattered signal was utilized

to view the crack propagation. For Al/SiC composites

the etching did not provide significant help. Therefore,both etched and unetched samples were tested.

A prenotch of about 1.01.5 mm in length was cut in-

side the specimens with the help of a 0.4 mm thick dia-

mond wheel. A special apparatus was then used to cut

very fine notches of about 1020 lm in radius with a

moving razor blade and very fine diamond paste. A dial-

gage indicator was integrated to the set-up to measure the

depth of the notch. Fig. 4 shows a single-edge-notched

specimen. The details of this set-up are provided by

Moon [16]. The prenotch and notch lengths and radii

were measured accurately with the help of an optical

microscope.

Table 2

Specimen details and fracture toughness calculations for composites

Sample used in

fracture tests

Width W

(mm)

Thickness B

(mm)

Notch length a

(mm)

Critical load (N) a=W Avg. KIC (MPaffiffiffiffi

mp

)

Cu/Al2O3 #1 4.95 4.26 1.865 200.0 0.376 4.0

Cu/Al2O3#2 5.47 4.71 1.774 267.0 0.324

Al/Al2O3 #1 5.56 6.26 2.194 502.0 0.394 6.5

Al/Al2O3 #2 5.4 5.94 1.933 600.0 0.358

Al/SiC #1 5.66 3.09 2.082 578.0 0.368 12.7

Al/SiC #2 5.72 3.06 1.5865 738.0 0.277



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4.1. Testing procedure

An environmental scanning electron microscope

(Electroscan model # 2020) was used to observe thein situ crack growth during the test. The microscope was

equipped with a loading stage. This equipment had a

long detector, which could scan the surface under low

vacuum. The load cell was aligned with the microscope

so that the electron beam could navigate and scan the

crack path as it propagated inside the sample. In situ

crack growth was recorded by an attached video cam-

era. Still pictures were obtained at various crack prop-

agation stages with the help of a digital imaging camera.

A vacuum of 2.55 Torr was maintained during the

fracture test to get clear images. The detector was placed

at the crack tip so that the crack propagation could be

followed and scanned by the electron microscope.

The notched specimens were placed in a three-point

bending configuration for testing. A pair of T shaped

steel fixtures with rollers were used to mount the speci-

men. The spacing between the rollers on the test fixture

was about 16 mm. The tails of these T fixtures were

gripped in between the flat grips of the loading cell.

A layer of soft acrylic tape padding was applied in be-

tween the specimen and rollers to prevent the grip-

movement during pumpdown. A schematic of this fix-

ture is shown in Fig. 5. Two tests were performed for

each composite to obtain consistent results. The loading

was applied under displacement control at the rate


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5.2. Al/Al2O3 composite

Fig. 8(a) shows a notched specimen before the

bending test. The micrograph shows the notch to be

surrounded by the metal phase. It was observed that

the crack did not initiate at the notch tip (Fig. 8(b)).

Instead, it initiated at the metalceramic interface and

then propagated inside the brittle ceramic. Further

propagation is shown in Fig. 8(c). The fracture mech-

anism inside this co-continuous composite was very

different from that in the Cu/Al2O3 composite. One

reason for crack propagation predominantly inside the

ceramic phase was the lower tensile thermal residual

stress inside the metal phase. The magnitude of tensile

stresses in the Al phase was only 112 MPa (due to a

large volume fraction of aluminum in the composite) as

compared to 620 MPa in the Cu phase of the Cu/Al2O3composite. Moreover, the ceramic grains were small

(0.31.0 lm in size) and the contiguity of ceramic

grains in this composite was considerably greater than

that in the Cu/Al2O3 composite (where each spherical

ceramic particle was a single grain). The large metal

spheres blunted the crack path and forced the crack to

grow around the sphere. This phenomenon is clearly

depicted in Fig. 8(d), which shows cracking of sample

#2. This mechanism led to significant toughening of

the composite. It is clear from an examination of the

fracture surface as shown in the SEM micrograph of

Fig. 9(a), that the metal phase experienced considerable

plastic deformation. There is evidence of void forma-

tion at the metalceramic interface, and at certain lo-

cations the voids grew interacting with the neighboring

voids, shown in Fig. 9(b). The ductile fracture at the

metal interface contributed to the fact that this com-

posite failed at considerably higher loads as compared

to the Cu/Al2O3 composite.

Fig. 6. Crack propagation in Cu/Al2O3 composite: (a) Cu/Al2O3 specimen before loading; (b) crack propagation along the metalceramic interface;

(c) further crack propagation along interface; (d) crack propagation inside the metal alloy.



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5.3. 6092 T6 Al/SiC composite

The crack initiation and growth at various stages are

shown in Figs. 10(a)(c). It was observed that several

small cracks initiated at the notch tip inside the metal

phase. A major crack emerged subsequently and grew

at 45 from the original notch. It grew in the metal

phase while the ceramic particles remained intact.There were several secondary branches close to the

crack tip, which grew at 45 from the main crack

(Fig. 10(c)). The fracture surface appeared to have

shear lip formations like those in unreinforced alumi-

num alloys, shown in Fig. 11. At higher magnifications,

small voids could be seen in the aluminum matrix close

to ceramic reinforcements. At higher magnifications,

small voids were seen even inside the metal phase, an

evidence of severe plastic deformation of matrix prior

to fracture. The specimens fractured at significantly

higher loads than those of the Cu/Al2O3 composite and

slightly higher than the Al/Al2O3 system. Although the

metal and ceramic volume fractions of this composite

were comparable to those of the Al/Al2O3 composite,

the crack growth behavior was very different. This is

essentially attributed to different microstructures and

degrees of contiguity of the ceramic phase in these two

composites. As the ceramic phase was in the form of

suspended particles, the crack path and fracture

mechanism was driven by characteristics of the ductile

matrix. An aluminum alloy specimen was fractured to

observe the crack propagation inside a pure metallic

alloy. It was found that the fracture characteristics

were similar to the composite. The following section

describes the fracture toughness computations for these

composites.

6. Fracture toughness calculations

In order to calculate fracture toughness, the load atwhich the crack initiated (and started to grow) at the

notch tip, was considered as the critical load. Linear

elastic fracture mechanics (LEFM) was used to estimate

the fracture toughness of the composites. In the three-

point bending configuration, the stress intensity factor

for a finite size sample is given by the expression de-

scribed in Anderson [17]:

KI PB

ffiffiffiffiffiW

p f aW

;

f aW

3S

WffiffiffiffiffiaW

r2 1 2 a

W

1 a

W

3=2

1:99 aW

1 aW

2:15 3:93 aW

2:7 a

W

2 : 1

In which, KI is the mode I stress intensity factor, P is the

applied load in the three-point bend configuration, a is

the crack length, W, B and S are the specimen width,

thickness and length, respectively. The critical stress

intensity factor KIC is defined at the critical load when

cracks initiate and grow at the notch tip. Table 2 lists the

stress intensity factors for the three composites for dif-ferent samples, calculated according to the above

equation. For Cu/Al2O3 composite, the KIC was calcu-

lated as 4.0 MPaffiffiffiffi

mp

. This is higher than that of

monolithic Al2O3 under bending. The KIC value for

monolithic Al2O3 ranges from 2.5 to 4.5 MPaffiffiffiffi

mp

de-

pending on the grain size [18].

The LEFM calculations gave an estimate for the

fracture toughness of Al/Al2O3 composite as 6.5 MPaffiffiffiffim

p. However, the metal phase in this composite has

significant plastic deformation and cavitation, therefore

these calculations may not be valid. The metalmatrix

composite 6092 Al/SiC had the maximum fracture

toughness value (12.7 MPaffiffiffiffi

mp ) by LEFM calculations.Even though the volume fraction of metal phase in both

Al/Al2O3 and Al/SiC composite were similar, the failure

mechanisms and fracture toughness values for both

composites were very different. This again confirms that

the fracture characteristics of metalceramic composites

are not a mere function of volume fraction; but micro-

structural properties like contiguity of phases and ther-

mal residual stresses play a significant role. The

following section describes finite element simulations to

model the effects of residual stresses in failure mecha-

nisms of co-continuous composites.

Fig. 7. Fractured surface of Cu/alumina composite showing no

breaking of ceramic particles.



8/12

7. Finite element simulations

In order to understand and explain the failure

mechanisms and experimental results, finite element

simulations were performed, using a special software,

Franc 2D. It was developed by Wawrzynek and

Ingraffea from Cornell University [11]. This software

enables adaptive remeshing during crack propagation.

Hence, it is possible to determine the stress distribution,

stress intensity factors and crack path during dynamic

crack propagation inside a structure.

A globallocal approach was utilized to simulate the

stresses and boundary conditions, applied during the

three-point bend tests. In the global model, the original

specimen dimensions were used and the stresses near the

notch tip were computed. The computed stresses were

used as boundary conditions in the local approach. For

the local model a small portion of the specimen near the

notch tip was used in order to simulate the microstruc-

ture as seen in the SEM and optical micrographs of the

composites. The residual stresses inside the metal and

ceramic phases were simulated and crack propagation

was studied at the microstructural scale. Details of the

two models are described in the following two sections.

7.1. Global model

The finite element model utilized for global approach

is shown in Fig. 12. Six-node triangular elements were

chosen for meshing. A point force was applied at the

opposite edge of the notch. The magnitude of the ap-

plied force was chosen to be the value when cracks

started to propagate inside the specimen. A simple roller

boundary condition was applied at the edge containing

the notch. The effective composite properties obtained

by micromechanical modeling were used for stiffness,

Fig. 8. Crack propagation in Al/Al2O3 composite: (a) notched Al/Al2O3 sample #1 before test; (b) crack propagation in sample #1; (c) crack

propagation in the brittle phase and along the interface; (d) sample #2, Al spheres, blocking the crack path.



9/12

Poissons ratio and density values of the composites in

the model. These values were obtained by using uniaxial

compression tests as well as micromechanical modeling.

Details are provided by Agrawal [13]. Table 1 lists these

values for the two composites.

Stress intensity factors were obtained by introducing

very small cracks at the tip of the Notch. Due to the

relatively large length of the notch, these values were

approximately constant over the length of small cracks.

Mode I was found to be the dominant fracture mode

and mode II was relatively insignificant. Average KIC

values of the composites were used to calculate thestresses around the notch tip. LEFM assumptions were

used throughout this modeling. The following expres-

sions from LEFM [17] were used for the stresses around

the notch tip:

rxx KIffiffiffiffiffiffiffi2pr

p cos h2

1 sin h

2

sin

3h

2

;

ryy KIffiffiffiffiffiffiffi2pr

p cos h2

1 sin h

2

sin

3h

2

;

rxy KI

ffiffiffiffiffiffiffi2prp cos h

2

sin

h

2

sin

3h

2

:

2

The stress distribution around the notch tip was calcu-

lated for both composites. These values were used as

boundary conditions for the local model.

7.2. Local model

In order to simulate crack propagation at the mi-crostructural level, a small portion of the composite near

the tip of the notch was designed to be similar to the

micrographs of the composite. This microstructural re-

gion was surrounded by effective homogeneous material,

shown in Fig. 13. The mechanical and fracture proper-

ties of the effective media as well as individual metal and

ceramic phases are listed in Table 1. A hexagonal

structure was used in order to simulate the spherical

grain shape in 2D. In order to simulate the grain size of

510 lm for Cu/Al2O3 and 50100 lm for Al/Al2O3,

some conversions had to be made for length units for

physical and mechanical properties of the metal and

ceramic phases inside the computational cell. Two types

of loads, thermal load due to thermal residual stresses

and mechanical stresses generated by three-point bend

tests were applied on this computational cell. The ef-

fective temperature DT (listed in Table 1) was used for

thermal residual stresses. The effective temperature was

chosen such that it would give rise to similar thermal

stress distribution inside metal and ceramic phases, as

measured by neutron diffraction experiment. In order to

apply bending stresses in the local model, the stress

distribution around the notch generated by the three-

point bend tests was used as boundary tractions. To

avoid rigid body rotation, x and y coordinates werefixed on one of the nodes. A small crack was introduced

at the notch tip and its growth was observed. The phase

surrounding the notch tip was alternatively taken as

metal and ceramic. Crack propagation was observed in

both situations. In order to observe the effects of ther-

mal and mechanical (bending) stresses, crack propaga-

tion was observed under three different conditions:

1. Pure thermal loading.

2. Pure mechanical loading.

3. Combined thermal and mechanical loading.

The growth was stopped as soon as the crack reached an

interface because the software did not have the capa-

bility to handle the interfacial fracture.

7.3. Cu/Al2O3 composite

This composite had about 70% of spherical Al2O3ceramic phase by volume. The light hexagonal phase

in the microstructural region in Fig. 13 represents the

ceramic phase. First, the crack tip was introduced in

the metallic region. The simulations were performed

for pure thermal and mechanical loading as well as

combined loading. Thermal and bending stresses were

Fig. 9. (a) Fracture surface of Al/Al2O3 (SEM micrograph) composite

showing yielding and cavitation in Al spheres. (b) Optical micrograph

showing the fractured parts of Al/Al2O3 composite.



10/12

linearly superimposed for the case of combined load-ing. The simulations confirmed that the stress distri-

bution under combined loading is still dominated by

the influence of thermal stresses.

A precrack was introduced at the tip and the value ofstress intensity factors, KI and KII, were obtained for

different loading conditions. The crack propagation was

observed under all three loading conditions mentioned

in the previous section. The maximum strain energy

criterion was used to compute the direction of crack

propagation. The crack did not grow in the presence of

pure thermal or bending stresses. However, in the case

of combined loading under the influence of combined

loading, the crack began to propagate in the direction of

high tensile stresses and stopped at an interface where

the stress discontinuity was fairly high as shown in Figs.

14(a)(c). Due to limitations of the finite element soft-

Fig. 10. Crack propagation in SiC/6092 T6 Al/SiC composite: (a) crack initiation in Al/SiC composite (sample #1) in the metal phase; (b) branching

and turning of crack (sample #1); (c) further propagation and branching (sample # 1); (d) zigzag crack path and yielding of matrix (sample # 2).

Fig. 11. Shear lip formation in Al/SiC composite. Fig. 12. Three-point bend specimen used for global analysis.



11/12

ware the simulations could not be performed for crack

growth near or at the interfacial zone.

The stress discontinuity and high tensile stresses close

to an interface made cracks propagate either along an

interface or inside the metallic phase for Cu/Al2O3composites. These simulations confirm the behavior

observed during the bend tests inside the ESEM cham-

ber as shown in Figs. 6(a)(d).

In order to investigate if the crack would grow in theceramic phase and observe the crack path, the micro-

structure was changed so that the elements close to the

notch tip were converted to ceramic phase. The crack

did not grow under combined loading. The simulations

confirmed the fact that the crack grows along the in-

terface or inside the metal phase in Cu/Al2O3 composite,

due to the presence of thermal stresses. The high thermal

residual stresses in the metallic phase due to its lower

volume fraction for this composite led to this unique

behavior.

7.4. Al/Al2O3 composite

The microstructure of this composite was compli-

mentary to the Cu/Al2

O3

composite. As opposed to Cu/

Al2O3, this composite had Al metal phase in the form of

spheres. The volume fraction of the spherical metal

phase was about 70%. Therefore, the model-design, that

was used for Cu/Al2O3 was utilized in this case as well,

except that the hexagonal phase was taken as metal and

the remainder was considered as ceramic. The size of the

metalspheres in this composite was in the range of 50

200 lm compared to 5 lm grain size of ceramic spheres

in Cu/Al2O3 composite. Therefore, a different scale of

local model was used for simulations.

First, the notch was extended inside the ceramic

region. The simulations were performed for all the

three cases, pure thermal, pure bending and combinedloading. By examining the figures, it can be concluded

that bending stresses play a more dominant role for

this composite. Thermal residual stresses caused a

compressive zone below the notch tip. The compres-

sive stress was about 90120 MPa. However, the

bending stresses gave rise to very high tensile stresses

that were an order of magnitude higher than com-

pressive stresses.

A small crack was introduced at the notch tip. The

simulations were performed to find the stress intensity

factors. There was a negative KI value due to compres-

sive loads inside the ceramic phase. However, this valuewas very small compared to high positive value of KIdue to bending stresses. Hence, the combined KI ex-

ceeded the critical value and the crack propagated inside

the ceramic phase. In order to see the individual con-

tributions of mechanical and thermal stresses, the sim-

ulations were also performed without residual stresses.

The crack path for the case of pure bending stresses was

very similar to the case for combined loading.

Fig. 14. Crack propagation inside Cu/Al2O3 composite under combined thermal and mechanical loading. (a) initial crack; (b) propagation after

3 steps of 1 lm increment and (c) after 5 lm increment.

Fig. 13. Material distribution for the local analysis.



12/12

In order to investigate if the crack would grow inside

the metallic phase, the tip elements were changed to

aluminum and simulations were performed for the same

loading conditions. The crack tip was found to be under

tension due to both thermal and mechanical stresses.

The tensile stress due to thermal residual stresses was

found to be significantly lower than that generated bymechanical stresses. The superimposed KI value was

found to be lower than the KIC value for aluminum

metal and the crack did not grow inside the metallic

phase. This confirms the behavior observed during ex-

periments where crack propagated around the metal

sphere and grew inside the ceramic phase.

8. Conclusions

The fracture characteristics of co-continuous metal

ceramic composites with different microstructures were

investigated and compared with metalmatrix compos-

ites. Neutron diffraction experiments were utilized to

measure the thermal residual stresses. The ceramic phase

in both Cu/Al2O3 and Al/Al2O3 composites was found

to be under compression and in both cases the metal

phase was found to be in tension. However, tensile stress

in the Cu phase was found to be significantly higher

compared to Al phase. This was due to differences in

volume fraction, higher melting point of copper alloy

and the different contiguity of composites.

Three-point bend tests were performed inside the

scanning electron microscope in order to observe crackpropagation at the microstructural level. It was found

that high tensile stresses in Cu phase, and at the Cu-

Al2O3 interfaces lead to crack propagation inside Cu

phase and at the interface. Whereas, for Al/Al3O3composite the crack propagated inside the ceramic

phase. In contrast, the fracture characteristics of metal

matrix composite (Al/SiC composite) were dominated

by the metal matrix.

The experimentally observed results for co-continu-

ous composites were also confirmed by finite element

simulations. The simulations provided the stress inten-

sity factors and stress distribution close to the notch tip

for various cases and helped in explaining the behavior

and influence of thermal residual stresses. They indi-

cated that presence of high thermal residual stresses

could influence the fracture behavior and change the

failure mechanism, as in the case of Cu/Al2O3 compos-

ite. In the absence of thermal residual stresses, the crack

path is governed by the individual fracture toughness of

the metal and ceramic phases and the plastic behavior of

the metallic phase. It can be concluded from this re-

search that contiguity and thermal residual stresses play

a very significant role in mechanical characteristics of

metalceramic composites.

Acknowledgements

This work was supported by an Army Research Of-

fice MURI Grant No. DAAH06-96-1-0331 to Purdue

University. Authors would also like to thank Dr. Frank

Cichocki for help in processing the composites and

Dr. Keith Bowman of Materials Science for his valuableadvice.

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