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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
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* 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.
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doi:10.1016/j.compscitech.2003.09.026
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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.
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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.
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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.
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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.
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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.
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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.
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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.
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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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