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alumina composites with interpenetrating networks
MATERIALS SCIENCE & ENGINEERING ELSEVIER
Materials Science and Engineering A197 (1995) 19 30
A
Strength and fracture toughness of aluminum/alumina composites with interpenetrating networks Helge Prielipp a, Mathias Knechtel a, Nils Claussen ", S.K. Streiffer b, H. Miillejans b, M. Rfihle b, Jfirgen R6deD* ~'Advanced Ceramics Group, Technische Universitdt Hamburg-Harburg, D-21073 Hamburg, Germany bMax- Plam'k-htstitut ./~ir MetallJbrschung, Institut fiir Werkstoffivissenschq#, D-70174 Stuttgart, Germa*o, Received 18 July 1994; in revised form 27 October 1994
Abstract
The mechanical properties of metal reinforced ceramics, especially AI/A1203 composites with interpenetrating networks, are described. Key parameters to tailor the characteristics of these materials are the ligament diameter and volume fraction of ductile reinforcement. Fracture strength and fracture toughness data are given as a function of both variables and are compared with the corresponding values for the porous preforms. A simple model accounts for the influence of metal volume and metal ligament diameter on the plateau toughness of the composites. The increase in fracture strength from the porous preform to the composite is found to be much larger than the gain which can be predicted from the increase in fracture toughness alone. A discussion of fracture strength in these composites therefore must include at least two issues, crack propagation through the matrix as well as crack initiation at metal filled pores.
Keywords: Fracture;
AI/A120 3 composites; Interpenetrating networks
1. Introduction
Metal reinforced ceramics have very attractive mechanical properties when designed as interpenetrating networks. Various methods are available to produce these materials, such as squeeze casting [1], directed metal oxidation [2], and infiltration with [3] and without [4] gas pressure. Potential advantages of these composites are high toughness and strength compared with the untoughened matrix material. A1/A1203 produced by directed metal oxidation [5] was reported to exhibit a fracture strength st- of 345 MPa and a fracture toughness Kit of 9.5 M P a m 12. For A1/AI:O 3 manufactured by gas pressure metal infiltration, a strength of 760 M P a and toughness of 5.8 M P a m 1:2 were observed [6]. A fracture toughness as high as 15 M P a m ~:2 has been reported in the W C / C o system [7]. A substantial amount of work, in the fields of both ceramics [8 12] and intermetallics [13 16] contributes *Present address: TH Darmstadt, Ceramics Group, D-4295 Darmstadt, Germany. Elsevier Science S.A. SSD1 0921-5093(94)09771-4
to our current understanding. The toughening mechanism believed to be effective in ceramic/metal composites is the plastic stretching of metallic inclusions bridging the advancing crack. The bridging ligaments exert closure stresses which reduce the stress intensity at the crack tip [17]. There is general agreement that the toughening c o n t r i b u t i o n by plastic deformation of the ductile phases is governed by the yield strength of the metal constituent and its uniaxial flow stress as established under the constraint of a more or less well bonded interface between metal and brittle matrix. The change in uniaxial stress p in the bridging ligament with crack opening 2u is represented by a stress displacement function p(u) which uniquely describes the reinforcement characteristic [18]. The shape of the p(u) function determines the way in which the crack resistance develops as the cracks grows; hence it controls the final steady state toughness. Experimental efforts at this stage are mainly concerned with the issue of crack propagation. These studies are aimed at determining the p(u) function in model
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H. Prielipp et al. / Materials Science and Engineerblg A 197 (1995) 19 30
systems [10,13,14,17] or are restricted to post-fracture investigations of the final ductile phase elongation [9,19]. In contrast, the strengthening mechanisms afforded by the inclusion of the ductile phase have not yet been explored in any detail. Pickard et al. [20] modeled the strength of A1/A1203 composites reinforced by SiC particles. The strength limiting defect was quantified by the SiC particle size with no consideration of the metal phase. Initial computations of the strength of ductile particle reinforced brittle matrix composites were presented by Bao and Zok [21]. The calculations were based on the p(u) function and the assumption that the initial flaw size Co and the elastic properties of the matrix are the same as those of the composite. The results showed an enhancement of toughness as well as strength, while the toughness increases more than the strength. Furthermore, the steady state toughness increases monotonically with debond length, whereas the strength is maximized at an intermediate value of debond length. A correlation between the calculations and experimental data has not yet been presented. In contrast to this theoretical work, a reduction in the initial flaw size in the composite was also proposed [6]. Factors which influence the toughening imparted by a ductile phase are the volume content, ligament diameter, interface properties and the metal properties such as flow stress, work hardening and ductility. Theoretical work [7,8,12,22] suggests that the fracture resistance will increase with volume fraction and bridge diameter. Good experimental data are so far restricted to the field of hard metals [7]. The intent of this paper is to provide a perspective of the range of mechanical properties which can be achieved with metal reinforced ceramics. Two key parameters in tailoring the characteristics of A1/A1203 composites are investigated, the metal volume fraction and metal ligament diameter. Metal infiltration affords an opportunity to study the influence of these parameters on the strength and toughness. Finally, this manufacturing method also makes it possible to examine the influence of metal properties and interface behavior on these same properties.
2. Experimental approach
starting with 55 vol.% A1 and 45 vol.% A1203 (Ceralox HPA 0.5). After attrition milling in acetone with alumina balls of 3 mm diameter for 9 h, the sieved and dried powder mixture was uniaxially pressed into plates at 5 0 M P a and then cold isostatically pressed at 250 MPa. The plates were reaction bonded with the following heating schedule: 7 h to 450 °C, then 20 h to 1150°C and a dwell time of 6h. Sintering at 1275 °C for 30min yielded a porous alumina body with a density of 75% theoretical density (TD). Medium- and coarse-sized microstructures resulted from slip cast medium (Alcoa CT 2000 SG) and coarse (Alcoa CL 5000) grained alumina powders respectively. Slurries with 45 vol.% solid content were cast into plates of dimensions 50 x 30 x 10ram 3. In contrast to the medium-grained alumina powder, the coarse powder was first attrition milled in ethanol with alumina balls of 3 mm diameter for 4h, and a gummi arabicum binder was added to the slurry. Sintering for l h at 1450 °C (medium-grained powder) and 1650 °C (coarsegrained powder) yielded bodies with a density of 75% TD. The porous plates were ground to dimensions of 50 x 30 x 5 m m 3 and then infiltrated with pure A1 (99.999%) in a specially designed gas pressure metal infiltrated furnace with an incorporated hydraulic ram (Fig. 1). The plates were originally held in a fixture which was immersed in a crucible filled with metal chips. The furnace was heated past the melting point of A1 (671 °C) up to 1050 °C in vacuum, and an argon pressure of 15 MPa was applied for 30 min to facilitate infiltration (Fig. l(a)). Subsequently, the furnace was cooled and, at 700 °C, the infiltrated plates were lifted out of the melt (Fig. l(b)). The pressure was not released before the temperature decreased below the melting point in order to suppress leakage of liquid metal from the preform. The metal infiltrated medium scale (m) and coarse scale (c) composites are compared with porous and dense A1203 of about equal grain size.
sampleholder crucible
I
ceramic body
The A1/A1203 composites were prepared by gas pressure metal infiltration. Three materials of different microstructural scale (termed small (s), medium (m) and coarse (c)) were produced with a metal content of 25 vol.%. Materials with medium-scale microstructure and varying metal content from 10 to 40 vol.% A1 were also produced. An additional batch of the coarsegrained material was prepared with 35 vol.% metal. A1203 bodies with small pore channels were manufactured using a reaction bonded alumina (RBAO) [23],
fullyinfiltrated
uninfiltrated
moltenmetal
infiltrated
a)
graphite-heater
b)
crucible support
Fig. 1. Schematic diagram of gas pressure metal infiltration furnace during infiltration (a) and after infiltration (b).
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H. Prielipp et al. / Materials Science and Engineering A 197 (1995) 19 30
The composites with medium microstructural scale are compared with a fine-grained alumina (Taimei T M DAR). This A1203 powder was uniaxially and subsequently cold isostatically pressed into bars. Sintering at 1350°C for 15rain yielded bodies with a density of greater than 99% TD. The manufacturing process is described in more detail elsewhere [24]. Sintering of the slip cast medium-grained A1203 powder (CT 2000 SG) compacts at 1650°C for 6 h yielded coarse-grained alumina with a density of greater than 98% TD. This A1203 was compared with the c composites. Densities were measured geometrically and by the Archimedes method using water as the immersion medium. The pore size distributions were determined by mercury porosimetry (poresizer 9320, Micrometrics). Microstructures of the porous and dense alumina were revealed by thermal etching in air at 1300 °C for 15 min and 1340 °C for 20 rain respectively. The grain sizes were determined by a linear intercept technique [25]. Optical microscopy of polished surfaces was used to characterize the microstructures. Scanning electron microscopy (SEM) was employed to investigate both fracture surfaces and surfaces which were polished and thermally etched. S, m and c composites with 25 vol.°/o metal content were further characterized by conventional transmission electron microscopy (TEM) using a J E O L 2000FX operated at 200 kV. T E M samples were prepared by standard dimpling and Ar ion-milling techniques. Bonding at AI-A1203 m e t a l - c e r a m i c interfaces in all three scale materials with 25 vol.% metal content was investigated by spatially resolved electron energyloss spectroscopy (EELS) in a dedicated scanning transmission electron microscope (STEM). This instrument has a beam diameter of less than 1 nm and an energy resolution better than 0.8 eV which allows investigation of near-edge structures in spectra with high spatial resolution (Vacuum Generators, V G HB501). Three spectra were recorded for each interface: one at the interface, one in the nearby metal and one in the adjacent alumina grains. The spectrum measured at the interface also contains components from both materials owing to the size of the electron beam. These components can be removed by subtracting the two reference spectra. The remaining component of the spectra represents the bonding of the atoms directly at the interface. Further details of this spatial difference method as applied to interfaces are provided by Bruley [26], Bruley et al. [27] and Mfillejans and Bruley [28]. Elastic properties of metal/ceramic composites with 25 vol.% and 35 vol.% AI were determined using a non-destructive ultrasonic technique [29]. Measurements were performed using plates with a thickness not less than 5 mm. The fracture strength was measured in four-point bending according to G e r m a n standard D I N 51110 but
1oo 80
s
60
o- 40 [J_
20 0 10-2
10-1
10 0
10-2
Pore size [pm]
Fig. 2. Intrusion pore size distributions of small (s, density 75 vol.%), medium (m, density 60 90 vol.%) and coarse (c, density 65 and 75 vol.'7,,) grained alumina. with rectangular bars cut from the plates, of reduced length (25 x 4 x 3 m m 3) and with loading spans of 10 and 20 mm. The tensile side was polished to a 3 ~tm surface finish and the edges were beveled. The fracture toughness was determined using the single edge precracked beam (SEPB) method ( D | N 51109) using five bars of the same size as described above. The precracks were 1 2 m m long. Samples were renotched before testing to leave metal-ligament bridging lengths Cu of less than 25% of the total (including the notched) crack size c. The selected normalized bridging lengths Cb/C were in the range 0.2 0.25 and the normalized crack lengths, were in the range 0.25 0.5. According to Zok and H o l m [30], the measured toughness is thereby overestimated by 10%-30%.
3. Results 3.1. M i c r o s t r u c t u r e o f porous prefi)rms
All three A1203 types with a porosity of 25% exhibited a rather narrow distribution of the intrusion channel size. The corresponding values for the median intrusion channel size were 0.08 lam, 0.25 lam and 0.8 ~tm for s, m and c materials with a porosity of 25% respectively. With varying porosity, the median pore size also changed but still remained in a rather narrow distribution band. Fig. 2 shows bands of pore sizes for two extreme density values for m and c materials compared with the pore size distribution of the s material with 25% porosity. The pore size distributions are not overlapping as revealed by median values between 0.1 and 0.25 ~tm for the m and between 0.8 and 1 lam for the c samples. For the medium-grained alumina in the range 2 5 % - 4 0 % porosity, the median channel diameter is nearly constant (d = 0.25 ~tm). With decreasing porosity from 25% to 10% and 35% to 25%, the pore size will decrease for the m alumina from 0.25 to 0.1 gm and for the c alumina from 1 to 0.8 ~tm respectively.
22
H. Prielipp et al.
Materials Science and Engineering A 197 (1995) 19 30
The grain sizes for the porous alumina bodies prepared with the medium- and coarse-grained powder were 1.2 and 4.0 ~tm respectively. This holds for densities ranging from 60% to 90°/,, TD. The fine-grained dense A1203 (Taimei T M - D A R ) exhibited an average grain size of 1.7 mm which is nearly the same as that of porous m alumina. Sintering of the slip cast medium-sized AI203 powder yielded an average grain size of 1.2 ~tm in the porous and 3.8 ~tm in the dense (greater than 98%) alumina. This is comparable with the grain size of the porous A1203 bodies prepared with the coarse-grained A1203 powder.
3.2. Effect of metal ligament diameter Representative micrographs of the m and c A1/AI 203 composites are provided in Figs. 3(a) and 3(b) respectively. The pores were nearly fully infiltrated so that the composite bodies had a porosity level of about 1%. The microstructures display good homogeneity and are consistent with the results of the narrow intrusion pore size distribution. The intrusion pore sizes used to describe
Fig. 3. Optical micrographs of medium-grained (a) and coarsegrained (b) AI/AI203 composites with 25 vol.% A1. The bright phase is the metal, the dark phase the alumina grains.
the three porous Al203 bodies represent bottlenecks. Average ligament diameters as measured from polished sections give values which are about a factor of 5 higher than the intrusion pore sizes. Nevertheless, irrespective of the parameter used to characterize the scale, comparisons between different scale microstructures indicate a true scale invariance. TEM confirmed that the A1 Al203 interfaces were well bonded in all three samples, and no interfacial failures were observed in the approximately 30 50 metal-filled pores examined in each specimen. This is in contrast to the observations in C u / A I 2 0 3 composites prepared by gas pressure metal infiltration where microcracking was observed at about 50% of interfaces in the large-scale composite, but was not detected in the smaller scale composites [31]. Representative bright field micrographs are shown in Fig. 4(a)-(c). The pores of the matrix are usually filled with single-crystal A1. Microfaceting of the A1203 grains was found in all of these samples, as is shown in Fig. 5. EELS spectra for all three specimens investigated were identical within statistical accuracy. Typical spectra are shown in Fig. 6. The presence of the interface component and the faceting (TEM) indicate that new interfaces have formed, which suggests strong bonding between metal and ceramic. The structure of the interface component is being investigated further and compared with calculated spectra based on simple structural interface models. Observations of fracture surfaces reveal scale invarient metal deformation (Fig. 7(a) and (b)). The metal ligaments neck to a point or ridge and cave out near the interface. Little debonding is visible at AI A1203 interfaces. In contrast to the monolithic alumina, essentially transgranular fracture of the A1203 grains can be observed. Fracture strength and fracture toughness (here defined and reported as plateau toughness) values of porous preforms as well as of the AI/AI2 03 composites containing 25 vol.% metal are given in Fig. 8(a) and (b). AI infiltration increases the fracture strength from values between 130 and 150 MPa for the porous preforms to values from 510 to 710 MPa for the composites. The maximum fracture strength occurs for the composite with medium ligament diameter (Fig. 8(a)). The fracture toughness (Fig. 8(b)) increases from 1.5 1 . 9 M P a m 1/2 for the porous alumina to 2.9 7.4 MPa m ~/2 for the metal-reinforced ceramics. Furthermore, the fracture toughness of the composites increases with increasing metal ligament diameter. In stress intensity factor notation, the equilibrium crack configuration with an applied stress intensity factor KA, a crack length c, and dependent fracture toughness KR(C) can be written as
23
H. PHelipp et ell. /Materials Science and Engineering A 197 (1995) 19 30
Fig. 5. TEM image of AI20 ~ microledges at the A1 A120 ~ interface (medium-scale microstructure, 25 vol.% metal).
ate for the respective crack configuration expressed as
K~,i(c) can
g(c,r)pi(r)dr
K~,i(c) =
be
(3)
While the underlying description in Eq. (3) appears very informative, it is not fundamental. The closure stresses p(r) are a function of the crack opening displacement (COD) 2u and will therefore change with crack length (which will affect the local COD). The notation of mechanical energy release rate G which implicitly includes closure stresses as a function of COD, may therefore be more applicable. The crack resistance term R(c) is written as the sum of a crack tip resistance term Ro and microstructural terms R~,,(c) which, again in equilibrium, is balanced by GA: GA = R0 q- Z R , , i ( c ) = R ( c )
(4)
i
I
'
I
'
I
'
i
5
I
,
Al203
4
~nterlaceAI/AI203
g °3 x
2
~'~
int~~ce~.~n~t x4
Fig. 4. Bright field TEM images of the (a) fine scale, (b) medium scale and (c) coarse scale microstructures containing 25 vol.% A1.
KA = Ko + Y' K~,, (c) = KR(c)
(2)
i"
The fi'acture toughness is seen as being composed of a crack tip toughness term Ko and microstructural terms K/,i(c), which sum up the closure stresses p,(r) of all reinforcements with a weight function g(c, r) appropri-
-0 .J 7o
8'o
9'0
1do
14 o
12o
Energy Loss (eV) Fig. 6. Typical electron energy-loss spectra of the AI L edge. By subtracting the two reference spectra measured in the metal and in the ceramic from the interface spectrum, the interface component is obtained which contains information about the bonding between metal and ceramic.
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H. Prielipp rt al. / Materials
Science and En#wering
Al97
(1995) 19-30
Combining Eqs. (4) and (6), the fracture toughness at the plateau of the R-curve KS,, given at a crack length c,, where the first active bridge fails, can be written as KS = ]Ki + EL RI,. matrix(c) + EL R,,, mrtaJ”*
(7)
where K,, = (R, EL)‘!* is the crack tip toughness of the composite. SEM fracture surface observations reveal that in all composites with metal content greater than 20 vol.%,, alumina grains fracture essentially in a transgranular manner. Transgranular as well as intergranular fracture is observed only with a metal fraction of 10 vol.%. Accordingly, the toughening contribution by grain bridging can be neglected, and Eq. (7) can be reduced to K, = [K:, + KR,,.
Fig. 7. Representative fracture surfaces coarse-grained (b) AI!A120, composites 2, Al.
.* s
of medium-grained (a) and with 25 vol.‘%, Al: I, AIZOq;
R,,,(c) results from the J-integral formalism [36] (Eq. (5)) where u* is the opening at the last active bridge: R,,, = 2j;
I” P,(u)du 0
The stress displacement function p(u) uniquely describes the reinforcement characteristics [18] and is responsible for the increasing fracture resistance with increasing crack length R(c). The toughening contribution by plastic deformation of the ductile phases is governed by the mechanical properties of the metal, the interface properties (debonding), the metal ligament diameter and its volume fraction j In the case of intercrystalline fracture of the matrix, interlocking grains will remain in contact across the crack faces and possibly form bridges providing additional crack closure stresses [37,38]. In the elastic case or elasticcplastic case under smallscale yielding conditions, we can use the relation between applied stress intensity factor KA and mechanical energy release rate G, (with EL Young’s modulus under plane strain conditions):
,,,cta,(~)l’~~
(9)
Furthermore, TEM and EELS investigations predict strong AlPA120, interfaces independent of metal ligament diameter. In addition, the metal reinforcements exhibit scale invariant plastic deformation. Under the assumption that the mechanical properties are also independent of ligament diameter, the toughening contribution of the ductile phase is only governed by geometry, i.e. volume fraction and ligament diameter given metal/ceramic system.
800 600 400 200 0 Porosity 25 vol-%
Al 25 vol-%
Porosity 25 vol-%
Al 25 ~01-96
Fracture strength (a) and fracture toughness (b) of AI/AI,O, composites with 25 vol.% metal and varymg hgament diameter compared with the corresponding values for the porous preforms.
H. Prielipp et al..' Materials Science and Engbteerh~g, A 197 (1995) 19 30
25
compared with an alumina prepared from fine-grained A1203 powder of similar medium grain size. The fracture strength decreases from 560 to 100MPa with decreasing porosity (Fig. l l(a)) following an exponential relationship according to Rhyskewitsch [32]: c~t.= ao e x p ( - n P )
(1)
where n is an adjustable parameter, P is the fractional porosity and a0 is the fracture strength of the dense alumina. It is interesting to note that the fracture strength of A1 reinforced A12 03 is even higher than that of dense alumina. In the range 10 25 vol.% A1, the fracture strength reveals a constant value of about 700 MPa. With the metal fraction increasing from 25 to 40 vol.%, the strength increases approximately linearly to 800 MPa. The fracture toughness (Fig. l l(b)) decreases linearly with increasing porosity from 3.6 to 0.8 MPa m 1'2 for the alumina with 40% porosity. The fracture toughness increases monotonically with increasing metal content, up to 7 MPa m 1'2 at 40 vol.% A1. This is an improvement of nearly 100% compared with the dense alumina. c composites exhibit a similar trend, i.e. both the fracture strength and fracture toughness are enhanced (Fig. 12(a)) to 5 1 0 M P a (25vo1.% AI) and 650 MPa
Fig. 9. Optical micrographs of medium-grained AI/AI20 ~ composites with (a) 10 vol.% and (b) 40 vol.% AI.
3.3. Roh, o f metal content
Optical micrographs of alumina with medium ligament diameter infiltrated with 10 vol.% and 40 vol.% aluminum are shown in Fig. 9(a) and (b). The microstructures are homogeneous and it appears that a transition from islands of metal to islands of ceramic takes place (at least in a two-dimensional view). The average ligament diameter, as measured for polished sections, is also approximately a factor of 10 larger than the intrusion pore size (cf. Section 3.2). Fig. 10(a) and (b) compares fracture surfaces of the microstructures shown in Fig. 9. At low metal contents the ligaments deform by necking to a sharp point (separate ligaments), while at higher metal contents deformation to chisel edges (connected ligaments) takes place. For composites with metal content greater than 20 vol.%, fracture is essentially transgranular. With a metal content of 10 vol.% transgranular as well as intergranular fracture was observed (Fig. 10(a)). The results of fracture strength and fracture toughness of the medium-scale materials vs. porosity and volume fraction of A1 are shown in Fig. l l(a) and (b). In order to minimize the influence of grain size on fracture strength and toughness, the m composites are
Fig. 10. Fracture surfaces of medium-grained AI/AI203 composites with (a) 10 vol.% and (b) 40 vol.% AI.
26
H. Prielipp et al. /Materials Science and Engineering A 197 (1995) 19 30 1000
r
medium grained
r
0- porous
800 rl
:E
"r
800 ~.
[3 - i n f i R r a t e d
650-+ 6(
coarse grained ,
~ i ~,
600
t(-
400
.o,,o
....... "~. . . . . . ~ . . . . . . . . . .
200
-
O" = 560 exp(-4 P)
o
-'
/ dense
10 a)
20
40
30
a)
EL
medium grained [ O-
porous
o_
D - infiltrated t~
D-
[3
m
0
12 ¸ 10 / 8
AI content
10.5 +-0.7 Coarse grained 7.4±0.5
/
6
/
cm
2
/
~-
0
u~
~ i3!ii:ill
4.6t 0.5 i" /
4
iiiii!ii i;i~i~il;_ i
1.St 0.2
(,9 09
~-
Porosity
Porosity/AI content [vol-%]
8
~
7 25 vol-% 35 vol-%
25 vol-% 35 vol-%
............ "~"..... 5"...........
2
CO
i
0 o
b)
10
i
20
i
30
~:~:~:~:~:
::i:::i:::i:i dense
o
1.5t 0.1
25 vol-% 35 vol-%
25 vol-% 35 vol-%
i
40
Porosity/AI content [vol-%]
Fig. 11. Fracture strength (a) and fracture toughness (b) of AI/AI303 composites with medium ligament diameter and varying metal content compared with the corresponding values for the porous preforms.
(35volY0) and to 7 . 4 M P a m ~/2 (25vo1.% AI) and 10.5 MPa m 1/2 (35 vol.% ) (Fig. 12(b)). The strength and toughness of monolithic alumina with a grain size of 4 mm decreases with increasing porosity from 370 MPa for the dense alumina to 70 MPa for coarse-grained alumina with a porosity of 35% and from 4.6 to 1 . 5 M P a m ~/2 respectively. In comparison with the dense alumina, the A1 reinforcement yields an improvement in strength and toughness of more than 100%. Independent of metal ligament diameter, the Young modulus decreases from 260 GPa to 230 GPa for 25 and 35 vol.% AI respectively. Poisson's ratio remains roughly constant at 0.28. With E = 400 GPa for alumina, E = 70 GPa for A1 and a Poisson's ratio of 0.28 for both, the Young modulus can be evaluated for all the composites, using the arithmetic median of Paul's lower and upper Young's modulus of the composite. The computed data show a relatively good agreement with experimental measurements.
b)
Porosity
AI content
Fig. 12. Fracture strength (a) and fracture toughness (b) of A1/AI20~ composites with coarse ligament diameter and varying metal content compared with the corresponding values of the porous preforms.
as grains can reinforce a ceramic matrix by providing closure stresses in the crack wake shielding the crack tip from the applied stress. Process zone shielding [33] and crack deflection [34] are further possible mechanisms; however, they have only small toughening potential in these composites when compared with crack bridging [8,35].
4.2. Long crack toughness Observations on fracture surfaces of different scale microstructures point to scale invariant metal deforma-
O"
crack
4. Discussion
(I 4.1. Fracture mechanics Possible reinforcement mechanisms in metal ceramic composites are shown in Fig. 13. Ductile phases as well
~iiiY
Fig. 13. Schematic diagram showing possible crack bridging mechanisms in metal ceramic composites: bridging by ductile phases (1) and bridging by matrix grains (2).
H. Prielipp et al.
Materials Science and Engineering, A 197 (1995) 19 30
_rntai,,x[ crac~
coarse
medium
a)
C
R
P
.**/o.-- ........ ,."
m
f U
b)
c)
Fig. 14. Schematic diagram describing p(u) functions (b) and Rcurves (c) for metal reinforced ceramics with varying ligament diameter (a), based on the assumption of scale independent metal deformation.
27
metal is a linear function of volume fraction 1/1 as well as of metal ligament diameter d owing to the microstructural scale invariance of plastic deformation. This implies according to Eq. (9) that variations in fracture toughness with metal content should obey a square root dependence on Vf for a given microstructural scale of the metal phase. Fig. 15 demonstrates good agreement between theory and experiment, except for the coarse-grained composite with a metal content of 25 vol.%. In order to estimate R~...... ,,~ from Eq. (5), the p(u) function was assumed to be constant up to a m a x i m u m crack opening where the first ligament fails. According to literature results [10] this opening is about the same size as the ligament diameter, which in turn can be estimated from optical micrographs to be roughly a factor of 5 larger than the pore channel size measured with mercury porosimetry. Thus the solid lines of Fig. 15 correspond to ligament diameters of 0.4, 1.25 and 4 jam. Combining Eqs. (9) and (5) results in a closure stress for the AI of 400 MPa, which exceeds the yield stress of pure AI by an order of magnitude. The authors relate this fact to the constrained plastic deformation as well as the minute microstructural scale of the single-crystal ligaments.
4.3. Composite .fracture strength tion and debond lengths. Consequently, p(u) functions can be derived from one assumed function and transformed into another microstructural scale simply by stretching the crack opening parameter in accordance with the ligament diameter (Fig. 14). Large reinforcements will therefore lead to smaller crack closure stresses at given small COD, which will be maintained up to higher crack opening displacements. This will result in R-curves for large reinforcements which can be characterized by small initial slope with a high peak thoughness. Therefore, the long-crack fracture toughness (as measured by SEPB) will increase with ligament diameter. The SEPB method yields fracture toughness values for bridged cracks of lengths between 200 and 400 jam, which for our materials is near the plateau of the R-curve and can therefore be represented by Eq. (9). The crack tip toughness Ko can be determined by a procedure [18] which relies on a measurement of near crack tip opening displacements and use of the crack profile description of Barenblatt [39]. A crack tip toughness of 2 . 0 M P a m 12 for A1/A1203 composites produced by direct metal oxidation has been estimated [40] which can be compared with values of 2.0 MPa m ','2 for alumina [18]. Fig. 15 shows the results of the fitting procedure according to Eq. (9), taking Ko as 2.0 MPa m ~'2 for all composites and assuming that crack bridging is due to ductile reinforcements only (Section 4.2). The sole assumption required is that the fracture energy contribution R t..... lal of the ductile
The increase in fracture strength with metal content may seem surprising at first, since AI has a yield strength much below the strength of alumina, thus any rule of mixture does not apply. To understand the strength of metal/ceramic composites, one needs to appreciate that, in interconnecting networks, the ceramic provides the failure site at the largest flaw, and the metal provides the fracture toughness. In the following sections, qualitative descriptions are provided for trends of composite strength, with the assumption that the initial flaw size of the porous alumina is the same as that of the respective composite.
12 &- small I • - medium I m
8
¢/) ¢/)
6
0-
o~ c¢,=
4
c~ "~ o I-
2
~
-e ~ EqarSe
± ~
_
m
0 0
10
20
30
40
AI content [vol-%] Fig. 15. Influence of metal ligament diameter and volume fraction on steady state toughness of AI/AI20~ composites. Full lines are representations of Eq. (9).
28
H. Prielipp et al. / Materials Science and Engineering A 197 (1995) 19 30
Gm
Gs Gc
..... // //
G,R //
////
y('*
~ " "
............ /,.::~<:'.:~........................... a c
...../~, "/ ..,, . / ~ "
//
increased metal fraction
//].~#~"
.,"/ ,~! • ........~>..,' /
G, R Rm
/-~ /
.....' .";'::!';;;" ......., " "~.~:.......................................
/ " ,.,;.S,I ,,**
Rs
....::..::-IS .....
IL
C
C Fig. 16. Schematic diagram describing R-curves of metal reinforced ceramics with varying metal ligament diameter. The tangency condition of the applied energy release rate GA (straight, dashed lines) with the R-curves yields the respective instability points.
4.3.1. Effect of metal ligament diameter In describing the fracture strength, the actual slope of the R-curve, as well as the initial flaw size, are important factors. Inserting the values for toughness, as measured with SEPB, and strength in Griffith's equation and solving for the crack length yields almost identical flaw sizes for the s, m and c alumina with 25% porosity. In the schematic diagram in Fig. 16, identical initial flaw sizes are therefore used for all composites. The c composite exhibits a rather shallow R-curve so that crack instability occurs with a large degree of stable crack growth; hence in spite of a large plateau toughness the strength is rather low. The finegrained micostructure exhibits only a very small Rcurve height, so that the increase in strength is not as dramatic. Conversely, the m material appears to combine initial flaw size and slope of the R-curve most favorably and therefore has the highest fracture strength of 710 MPa. Due attention must be paid to the scale of the initial flaw size. In cases where this parameter is increased (as in thermal shock), the large-scale microstructure may prove to be advantageous compared with the fine- or medium-scale counterparts [41].
Fig. 17. Schematic diagram describing R-curves of metal reinforced ceramics with varying metal content. The tangency condition of the applied energy release rate GA (straight, dashed lines) with the R-curves yielded the respective instability points.
4.3.3. Effect of flaw bridging Our qualitative considerations describe only the trend of strength as a function of ligament diameter as well as of volume fraction. Assuming the same initial flaw sizes for the composites as for the corresponding porous alumina, the instability criterion (tangency condition of the applied stres intensity factor with Rcurve) predicts an enhancement of the composite strength directly proportional to the fracture toughness at the instability crack length. This statement is consistent with the initial modeling from Bao and Zok [21]. Quite on the contrary, however, our results show a larger enhancement of strength than of toughness owing to metal infiltration, as demonstrated in Fig. 18 by the mechanical property data normalized to the data of the porous preform. The above model assumes that the initial flaw is not bridged by the ductile phase. Based on investigations on the fracture strength of alumina [24], where the largest pore determines the failure origin, we assume that the site of the initial flaw remains unchanged but is converted to a metal filled cavity. Therefore the crack is initially bridged by metal with a volume fraction fb of 10
4.3.2. Effect of metal content A large metal content results from a preform with high porosity. In applying the Griffith equation, it appears that the initial flaw size increases with increasing porosity. Depending on the exact interaction of initial pore size and slope of the R-curve, an increase or decrease in strength with metal content may result. In the case of AI reinforcement, the Rcurves are steep enough to compensate for the small increase in initial flaw size and therefore yield an increase in strength with increasing metal content as shown schematically by the tangency condition in Fig. 17.
medium grained © - Strength [] - Toughness
~8
.o .. O' /' . .O" ....O ......
~(
~
2
,:::: ::::::?
El " "
. ..,,D ..
....
......
10
20
30
40
AI content [vol-%]
Fig. 18. Effect of AI content on the relative enhancement of the composite strength and toughness. The composite properties were normalized to the corresponding values for the porous alumina preforms.
29
H. Prielipp et al./ Materials Sciem'e and EnghwerhTg A 197 (1995) 19 30
long crack
metal
finitial : fb = 1 0 0 %
,/'
G, R
// a)
1
.
G,R
/,/'/ / c
b)
c
Fig. 20. R-curves (a,b) explain schematically the short crack anomaly (b) compared with long crack behavior (a).
fb = F(C) > f
1
small crack
fb = f = const.
C
Fig. 19. Hypothesis for crack initiation in metal/ceramic composites where failure-causing pores are eliminated and failure initiates at large metal filled cavities.
100% (Fig. 19(a)). As the crack extends into the matrix, it will consecutively encounter an increasing number of metal ligaments (Fig. 19(b)). If the crack length is large with respect to the metal filled cavity, the metal bridging fraction is constant (f, = J ) (Fig. 19(c)). Inclusion of the change in metal volume fraction leads to a modification of Eq. (5), where now f has to be replaced by
tion into tailored porous preforms. These materials exhibit fracture toughnesses up to 10.5 MPa m j'2 and fracture strength up to 810 MPa. In contrast to metal matrix composites (e.g. [42]), both the fracture strength and fracture toughness are increased with second phase content. TEM and EELS characterization show continuous, microcrack-free bonds in all composites of varying microstructural scale. Long-crack fracture toughness data can be inferred by using the plateau toughness of one sample microstructure and then scaling the toughness with respect to ligament diameter and metal volume, where the observation of scale invariant plastic deformation is utilized. Trends in fracture strength as a function of metal content and metal ligament diameter can be explained. A striking result, however, is the fact that the fracture strength increases more pronouncedly by metal infiltration than the plateau toughness derived from long crack measurements. This fact is attributed to large closure stresses during and after crack initiation at metal filled cavities associated with plastic deformation of these large metal spheres.
f(c): t¢i*
R,,= 2f(c)
Ij
pi(u)du
(10)
The foregoing discussions on the influence of the matrix on fracture are therefore still valid, but have to be augmented by the closure effect of the large metal filled pore. In addition, the initial slope of the R-curve is steeper than that of a long-crack R-curve (Fig. 20(a) and (b)). This qualitative description, by separating small- and long-crack fracture toughness, can explain the fact that the composite strength increases more than the (measured) long crack fracture toughness would predict (tangency condition in Fig. 20(b)).
Acknowledgements
We thank Maria Sycha for preparing the TEM specimens and Christina Scheu for performing part of the experiments and analysis. S.K.S. wishes to thank the Alexander von Humboldt-Stiftung for financial support. This work was supported by the Deutsche Forschungsgemeinschaft under contract number Ro 954/1-1 and the Volkswagen Foundation under contract number 1/66 760.
References 5. Conclusion
A1/AI203 composites can be produced with widely differing microstructures (especially metal volume and metal ligament diameter) by gas pressure metal infiltra-
[l] F. Lange, B.V. Velamakanni and A.G. Evans, J. Am. ('cram. Soe., 73 (2) (1990) 388 393. [2] M.S. Newkirk, A.W. Urquhart, H.R. Zwicker and E. Breval, J. Mater. Res., 1 (1) (1986) 81 89. [3] N. Travitzky and N. Claussen, J. Eur. Ceram. Soe.. 9 (1992) 61 65.
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H. Prielipp et al. / Materials Science and Engineering A 197 (1995) 19-30
[4] C. Toy and W.D. Scott, J. Am. Ceram. Sot'., 73 (1) (1990) 97-101. [5] M.K. Aghajanian, N.H. Macmillan, C.R. Kennedy, S.J. Luszcz and R. Roy, J. Mater. Sci., 24 (1989) 658-670. [6] S. Wu, A.J. Gesing, N.A. Travitzky and N. Claussen, J. Eur. Ceram. Sot'., 7 (1991) 277 281. [7] K.S. Ravichandran, Acta Metall. Mater., 42 (1) (1994) 143 150. [8] L.S. Sigl, P.A. Mataga, B.J. Dalgleish, R.M. McMeeking and A.G. Evans, Aeta Metall., 36 (4) (1988) 945-953. [9] B.D. Elinn, M. Riihle and A.G. Evans, Acta Metall., 37 (11) (1989) 3001 3006. [10] M.F. Ashby, F.J. Blunt and M. Bannister, Acta Metall. Mater., 37 (7) (1989) 1847 1857. [11] P.A. Mataga, Acta Metall., 37 (12) (1989) 3349-3359. [12] G. Bao and C.-Y. Hui, Int. J. Solid Struct., 26 (5 6) (1990) 631 642. [13] H.C. Cao, B.J. Dalgleish, H.E. Deve, C. Elliott, A.G. Evans, R. Mehrabian and G.R. Odette, Acta Metall., 37 (11) (1989) 2969 2977. [14] H.E. Deve, A.G. Evans, G.R. Odette, R. Mehrabian, M.L. Emiliani and R.J. Hecht, Acta Metall., 37(8) (1990) 1491 1502. [15] T.C. Lu, A.G. Evans, R.J. Hecht and R. Mehrabian, Ac'ta Metall. Mater., 39 (1991) 1853 1862. [16] L. Xiao and R. Abbaschian, Mater. Sci. Eng., A155 (1992) 135 145. [17] M. Bannister, H. Shercliff, G. Bao, F. Zok and M.F. Ashby, Acta Metall. Mater., 40 (7) (1992) 1531 1537. [18] J. R6del, J.F. Kelly and B.R. Lawn, J. Am. Ceram. Soc., 73(11) (1990) 3313 3318. [19] B.D. Flinn, C.S. Lo, F.W. Zok and A.G. Evans, J. Am. Ceram. Soc., 76 (2) (1993) 369 375. [20] S.M. Pickard, E. Manor, H. Ni, A.G. Evans and R. Mehrabian, Acta Metall. Mater., 40 (1) (1992) 177 184.
[21] G. Bao and F. Zok, Acta Metall. Mater., 41 (12) (1993) 35153524. [22] E.G. Evans and R.M. McMeeking, Acta Metall., 34 (12) (1986) 2435-2441. [23] S. Wu, D. Holz and N. Claussen, J. Am. Ceram. Soc., 76 (4) (1993) 970-980. [24] J. Seidel, N. Claussen and J. R6del, J. Eur. Ceram. Soc., in press. [25] J.C. Wurst and J.A. Nelson, J. Am. Ceram. Soe., 55 (2) (1972) 109. [26] J. Bruley, Microsc. Mieroanal. Microstruct., 4 (1993) 23. [27] J. Bruley, R. Byrdson, H. Mfillejans, J. Mayer, G. Gutekunst, W. Mader, D. Knauss and M. R/.ihle, submitted to J. Mater. Res.
[28] H. Mfillejans and J. Bruley, Ultramicroscopy, in press. [29] M.C. Bhardwaj, Am. Ceram. Soc. Bull., 69 (9) (1990) 1490. [30] F. Zok and C.L. Holm, Acta Metall. Mater., 38 (10) (1990) 1895 - 1904. [31] M. Knechtel, H. Prielipp, N. Claussen, S.K. Streiffer, C. Scheu, M. Rfihle and J. R6del, to be submitted to J. Am. Ceram. Soc. [32] Rhyskewitsch, J. Am. Ceram. Soc., 36 (2) (1953) 65-68. [33] A.G. Evans and R.M. Cannon, Acta Metall., 34 (5) (1986) 761 800. [34] K.T. Faber and A.G. Evans, Acta Metall., 31 (1983) 577-584. [35] J. R6del, J, Eur. Ceram. Soe., 10 (3) (1992) 143 150. [36] J.R. Rice, J. Appl. Mech., 35 (1968) 379-386. [37] R. Knehans and R.W. Steinbrech, J. Mater. Sci. Lett., 1 (1982) 327 329. [38] J. R6del, J. Eur. Ceram. Soc., 9 (4) (1992) 323 334. [39] G.I. Barenblatt, Adv. Appl. Mech., 7 (1962) 55-129. [40] J. R6del, M. Sindel, M. Dransmann, R.W. Steinbrech and N. Claussen, J. Eur. Ceram. Soc., 14 (1994) 153-161. [41] S. Sch6n, H. Prielipp, R. Janssen, J. R6del and N. Claussen, J. Am. Ceram. Soc., 77 (3) (1994) 701 704. [42] I.A. lbrahim, F.A. Mohamed and E.J. Laverna, J. Mater. Sci., 26(1991) 1137-1156.
Materials Science and Engineering A197 (1995) 19 30
A
Strength and fracture toughness of aluminum/alumina composites with interpenetrating networks Helge Prielipp a, Mathias Knechtel a, Nils Claussen ", S.K. Streiffer b, H. Miillejans b, M. Rfihle b, Jfirgen R6deD* ~'Advanced Ceramics Group, Technische Universitdt Hamburg-Harburg, D-21073 Hamburg, Germany bMax- Plam'k-htstitut ./~ir MetallJbrschung, Institut fiir Werkstoffivissenschq#, D-70174 Stuttgart, Germa*o, Received 18 July 1994; in revised form 27 October 1994
Abstract
The mechanical properties of metal reinforced ceramics, especially AI/A1203 composites with interpenetrating networks, are described. Key parameters to tailor the characteristics of these materials are the ligament diameter and volume fraction of ductile reinforcement. Fracture strength and fracture toughness data are given as a function of both variables and are compared with the corresponding values for the porous preforms. A simple model accounts for the influence of metal volume and metal ligament diameter on the plateau toughness of the composites. The increase in fracture strength from the porous preform to the composite is found to be much larger than the gain which can be predicted from the increase in fracture toughness alone. A discussion of fracture strength in these composites therefore must include at least two issues, crack propagation through the matrix as well as crack initiation at metal filled pores.
Keywords: Fracture;
AI/A120 3 composites; Interpenetrating networks
1. Introduction
Metal reinforced ceramics have very attractive mechanical properties when designed as interpenetrating networks. Various methods are available to produce these materials, such as squeeze casting [1], directed metal oxidation [2], and infiltration with [3] and without [4] gas pressure. Potential advantages of these composites are high toughness and strength compared with the untoughened matrix material. A1/A1203 produced by directed metal oxidation [5] was reported to exhibit a fracture strength st- of 345 MPa and a fracture toughness Kit of 9.5 M P a m 12. For A1/AI:O 3 manufactured by gas pressure metal infiltration, a strength of 760 M P a and toughness of 5.8 M P a m 1:2 were observed [6]. A fracture toughness as high as 15 M P a m ~:2 has been reported in the W C / C o system [7]. A substantial amount of work, in the fields of both ceramics [8 12] and intermetallics [13 16] contributes *Present address: TH Darmstadt, Ceramics Group, D-4295 Darmstadt, Germany. Elsevier Science S.A. SSD1 0921-5093(94)09771-4
to our current understanding. The toughening mechanism believed to be effective in ceramic/metal composites is the plastic stretching of metallic inclusions bridging the advancing crack. The bridging ligaments exert closure stresses which reduce the stress intensity at the crack tip [17]. There is general agreement that the toughening c o n t r i b u t i o n by plastic deformation of the ductile phases is governed by the yield strength of the metal constituent and its uniaxial flow stress as established under the constraint of a more or less well bonded interface between metal and brittle matrix. The change in uniaxial stress p in the bridging ligament with crack opening 2u is represented by a stress displacement function p(u) which uniquely describes the reinforcement characteristic [18]. The shape of the p(u) function determines the way in which the crack resistance develops as the cracks grows; hence it controls the final steady state toughness. Experimental efforts at this stage are mainly concerned with the issue of crack propagation. These studies are aimed at determining the p(u) function in model
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H. Prielipp et al. / Materials Science and Engineerblg A 197 (1995) 19 30
systems [10,13,14,17] or are restricted to post-fracture investigations of the final ductile phase elongation [9,19]. In contrast, the strengthening mechanisms afforded by the inclusion of the ductile phase have not yet been explored in any detail. Pickard et al. [20] modeled the strength of A1/A1203 composites reinforced by SiC particles. The strength limiting defect was quantified by the SiC particle size with no consideration of the metal phase. Initial computations of the strength of ductile particle reinforced brittle matrix composites were presented by Bao and Zok [21]. The calculations were based on the p(u) function and the assumption that the initial flaw size Co and the elastic properties of the matrix are the same as those of the composite. The results showed an enhancement of toughness as well as strength, while the toughness increases more than the strength. Furthermore, the steady state toughness increases monotonically with debond length, whereas the strength is maximized at an intermediate value of debond length. A correlation between the calculations and experimental data has not yet been presented. In contrast to this theoretical work, a reduction in the initial flaw size in the composite was also proposed [6]. Factors which influence the toughening imparted by a ductile phase are the volume content, ligament diameter, interface properties and the metal properties such as flow stress, work hardening and ductility. Theoretical work [7,8,12,22] suggests that the fracture resistance will increase with volume fraction and bridge diameter. Good experimental data are so far restricted to the field of hard metals [7]. The intent of this paper is to provide a perspective of the range of mechanical properties which can be achieved with metal reinforced ceramics. Two key parameters in tailoring the characteristics of A1/A1203 composites are investigated, the metal volume fraction and metal ligament diameter. Metal infiltration affords an opportunity to study the influence of these parameters on the strength and toughness. Finally, this manufacturing method also makes it possible to examine the influence of metal properties and interface behavior on these same properties.
2. Experimental approach
starting with 55 vol.% A1 and 45 vol.% A1203 (Ceralox HPA 0.5). After attrition milling in acetone with alumina balls of 3 mm diameter for 9 h, the sieved and dried powder mixture was uniaxially pressed into plates at 5 0 M P a and then cold isostatically pressed at 250 MPa. The plates were reaction bonded with the following heating schedule: 7 h to 450 °C, then 20 h to 1150°C and a dwell time of 6h. Sintering at 1275 °C for 30min yielded a porous alumina body with a density of 75% theoretical density (TD). Medium- and coarse-sized microstructures resulted from slip cast medium (Alcoa CT 2000 SG) and coarse (Alcoa CL 5000) grained alumina powders respectively. Slurries with 45 vol.% solid content were cast into plates of dimensions 50 x 30 x 10ram 3. In contrast to the medium-grained alumina powder, the coarse powder was first attrition milled in ethanol with alumina balls of 3 mm diameter for 4h, and a gummi arabicum binder was added to the slurry. Sintering for l h at 1450 °C (medium-grained powder) and 1650 °C (coarsegrained powder) yielded bodies with a density of 75% TD. The porous plates were ground to dimensions of 50 x 30 x 5 m m 3 and then infiltrated with pure A1 (99.999%) in a specially designed gas pressure metal infiltrated furnace with an incorporated hydraulic ram (Fig. 1). The plates were originally held in a fixture which was immersed in a crucible filled with metal chips. The furnace was heated past the melting point of A1 (671 °C) up to 1050 °C in vacuum, and an argon pressure of 15 MPa was applied for 30 min to facilitate infiltration (Fig. l(a)). Subsequently, the furnace was cooled and, at 700 °C, the infiltrated plates were lifted out of the melt (Fig. l(b)). The pressure was not released before the temperature decreased below the melting point in order to suppress leakage of liquid metal from the preform. The metal infiltrated medium scale (m) and coarse scale (c) composites are compared with porous and dense A1203 of about equal grain size.
sampleholder crucible
I
ceramic body
The A1/A1203 composites were prepared by gas pressure metal infiltration. Three materials of different microstructural scale (termed small (s), medium (m) and coarse (c)) were produced with a metal content of 25 vol.%. Materials with medium-scale microstructure and varying metal content from 10 to 40 vol.% A1 were also produced. An additional batch of the coarsegrained material was prepared with 35 vol.% metal. A1203 bodies with small pore channels were manufactured using a reaction bonded alumina (RBAO) [23],
fullyinfiltrated
uninfiltrated
moltenmetal
infiltrated
a)
graphite-heater
b)
crucible support
Fig. 1. Schematic diagram of gas pressure metal infiltration furnace during infiltration (a) and after infiltration (b).
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H. Prielipp et al. / Materials Science and Engineering A 197 (1995) 19 30
The composites with medium microstructural scale are compared with a fine-grained alumina (Taimei T M DAR). This A1203 powder was uniaxially and subsequently cold isostatically pressed into bars. Sintering at 1350°C for 15rain yielded bodies with a density of greater than 99% TD. The manufacturing process is described in more detail elsewhere [24]. Sintering of the slip cast medium-grained A1203 powder (CT 2000 SG) compacts at 1650°C for 6 h yielded coarse-grained alumina with a density of greater than 98% TD. This A1203 was compared with the c composites. Densities were measured geometrically and by the Archimedes method using water as the immersion medium. The pore size distributions were determined by mercury porosimetry (poresizer 9320, Micrometrics). Microstructures of the porous and dense alumina were revealed by thermal etching in air at 1300 °C for 15 min and 1340 °C for 20 rain respectively. The grain sizes were determined by a linear intercept technique [25]. Optical microscopy of polished surfaces was used to characterize the microstructures. Scanning electron microscopy (SEM) was employed to investigate both fracture surfaces and surfaces which were polished and thermally etched. S, m and c composites with 25 vol.°/o metal content were further characterized by conventional transmission electron microscopy (TEM) using a J E O L 2000FX operated at 200 kV. T E M samples were prepared by standard dimpling and Ar ion-milling techniques. Bonding at AI-A1203 m e t a l - c e r a m i c interfaces in all three scale materials with 25 vol.% metal content was investigated by spatially resolved electron energyloss spectroscopy (EELS) in a dedicated scanning transmission electron microscope (STEM). This instrument has a beam diameter of less than 1 nm and an energy resolution better than 0.8 eV which allows investigation of near-edge structures in spectra with high spatial resolution (Vacuum Generators, V G HB501). Three spectra were recorded for each interface: one at the interface, one in the nearby metal and one in the adjacent alumina grains. The spectrum measured at the interface also contains components from both materials owing to the size of the electron beam. These components can be removed by subtracting the two reference spectra. The remaining component of the spectra represents the bonding of the atoms directly at the interface. Further details of this spatial difference method as applied to interfaces are provided by Bruley [26], Bruley et al. [27] and Mfillejans and Bruley [28]. Elastic properties of metal/ceramic composites with 25 vol.% and 35 vol.% AI were determined using a non-destructive ultrasonic technique [29]. Measurements were performed using plates with a thickness not less than 5 mm. The fracture strength was measured in four-point bending according to G e r m a n standard D I N 51110 but
1oo 80
s
60
o- 40 [J_
20 0 10-2
10-1
10 0
10-2
Pore size [pm]
Fig. 2. Intrusion pore size distributions of small (s, density 75 vol.%), medium (m, density 60 90 vol.%) and coarse (c, density 65 and 75 vol.'7,,) grained alumina. with rectangular bars cut from the plates, of reduced length (25 x 4 x 3 m m 3) and with loading spans of 10 and 20 mm. The tensile side was polished to a 3 ~tm surface finish and the edges were beveled. The fracture toughness was determined using the single edge precracked beam (SEPB) method ( D | N 51109) using five bars of the same size as described above. The precracks were 1 2 m m long. Samples were renotched before testing to leave metal-ligament bridging lengths Cu of less than 25% of the total (including the notched) crack size c. The selected normalized bridging lengths Cb/C were in the range 0.2 0.25 and the normalized crack lengths, were in the range 0.25 0.5. According to Zok and H o l m [30], the measured toughness is thereby overestimated by 10%-30%.
3. Results 3.1. M i c r o s t r u c t u r e o f porous prefi)rms
All three A1203 types with a porosity of 25% exhibited a rather narrow distribution of the intrusion channel size. The corresponding values for the median intrusion channel size were 0.08 lam, 0.25 lam and 0.8 ~tm for s, m and c materials with a porosity of 25% respectively. With varying porosity, the median pore size also changed but still remained in a rather narrow distribution band. Fig. 2 shows bands of pore sizes for two extreme density values for m and c materials compared with the pore size distribution of the s material with 25% porosity. The pore size distributions are not overlapping as revealed by median values between 0.1 and 0.25 ~tm for the m and between 0.8 and 1 lam for the c samples. For the medium-grained alumina in the range 2 5 % - 4 0 % porosity, the median channel diameter is nearly constant (d = 0.25 ~tm). With decreasing porosity from 25% to 10% and 35% to 25%, the pore size will decrease for the m alumina from 0.25 to 0.1 gm and for the c alumina from 1 to 0.8 ~tm respectively.
22
H. Prielipp et al.
Materials Science and Engineering A 197 (1995) 19 30
The grain sizes for the porous alumina bodies prepared with the medium- and coarse-grained powder were 1.2 and 4.0 ~tm respectively. This holds for densities ranging from 60% to 90°/,, TD. The fine-grained dense A1203 (Taimei T M - D A R ) exhibited an average grain size of 1.7 mm which is nearly the same as that of porous m alumina. Sintering of the slip cast medium-sized AI203 powder yielded an average grain size of 1.2 ~tm in the porous and 3.8 ~tm in the dense (greater than 98%) alumina. This is comparable with the grain size of the porous A1203 bodies prepared with the coarse-grained A1203 powder.
3.2. Effect of metal ligament diameter Representative micrographs of the m and c A1/AI 203 composites are provided in Figs. 3(a) and 3(b) respectively. The pores were nearly fully infiltrated so that the composite bodies had a porosity level of about 1%. The microstructures display good homogeneity and are consistent with the results of the narrow intrusion pore size distribution. The intrusion pore sizes used to describe
Fig. 3. Optical micrographs of medium-grained (a) and coarsegrained (b) AI/AI203 composites with 25 vol.% A1. The bright phase is the metal, the dark phase the alumina grains.
the three porous Al203 bodies represent bottlenecks. Average ligament diameters as measured from polished sections give values which are about a factor of 5 higher than the intrusion pore sizes. Nevertheless, irrespective of the parameter used to characterize the scale, comparisons between different scale microstructures indicate a true scale invariance. TEM confirmed that the A1 Al203 interfaces were well bonded in all three samples, and no interfacial failures were observed in the approximately 30 50 metal-filled pores examined in each specimen. This is in contrast to the observations in C u / A I 2 0 3 composites prepared by gas pressure metal infiltration where microcracking was observed at about 50% of interfaces in the large-scale composite, but was not detected in the smaller scale composites [31]. Representative bright field micrographs are shown in Fig. 4(a)-(c). The pores of the matrix are usually filled with single-crystal A1. Microfaceting of the A1203 grains was found in all of these samples, as is shown in Fig. 5. EELS spectra for all three specimens investigated were identical within statistical accuracy. Typical spectra are shown in Fig. 6. The presence of the interface component and the faceting (TEM) indicate that new interfaces have formed, which suggests strong bonding between metal and ceramic. The structure of the interface component is being investigated further and compared with calculated spectra based on simple structural interface models. Observations of fracture surfaces reveal scale invarient metal deformation (Fig. 7(a) and (b)). The metal ligaments neck to a point or ridge and cave out near the interface. Little debonding is visible at AI A1203 interfaces. In contrast to the monolithic alumina, essentially transgranular fracture of the A1203 grains can be observed. Fracture strength and fracture toughness (here defined and reported as plateau toughness) values of porous preforms as well as of the AI/AI2 03 composites containing 25 vol.% metal are given in Fig. 8(a) and (b). AI infiltration increases the fracture strength from values between 130 and 150 MPa for the porous preforms to values from 510 to 710 MPa for the composites. The maximum fracture strength occurs for the composite with medium ligament diameter (Fig. 8(a)). The fracture toughness (Fig. 8(b)) increases from 1.5 1 . 9 M P a m 1/2 for the porous alumina to 2.9 7.4 MPa m ~/2 for the metal-reinforced ceramics. Furthermore, the fracture toughness of the composites increases with increasing metal ligament diameter. In stress intensity factor notation, the equilibrium crack configuration with an applied stress intensity factor KA, a crack length c, and dependent fracture toughness KR(C) can be written as
23
H. PHelipp et ell. /Materials Science and Engineering A 197 (1995) 19 30
Fig. 5. TEM image of AI20 ~ microledges at the A1 A120 ~ interface (medium-scale microstructure, 25 vol.% metal).
ate for the respective crack configuration expressed as
K~,i(c) can
g(c,r)pi(r)dr
K~,i(c) =
be
(3)
While the underlying description in Eq. (3) appears very informative, it is not fundamental. The closure stresses p(r) are a function of the crack opening displacement (COD) 2u and will therefore change with crack length (which will affect the local COD). The notation of mechanical energy release rate G which implicitly includes closure stresses as a function of COD, may therefore be more applicable. The crack resistance term R(c) is written as the sum of a crack tip resistance term Ro and microstructural terms R~,,(c) which, again in equilibrium, is balanced by GA: GA = R0 q- Z R , , i ( c ) = R ( c )
(4)
i
I
'
I
'
I
'
i
5
I
,
Al203
4
~nterlaceAI/AI203
g °3 x
2
~'~
int~~ce~.~n~t x4
Fig. 4. Bright field TEM images of the (a) fine scale, (b) medium scale and (c) coarse scale microstructures containing 25 vol.% A1.
KA = Ko + Y' K~,, (c) = KR(c)
(2)
i"
The fi'acture toughness is seen as being composed of a crack tip toughness term Ko and microstructural terms K/,i(c), which sum up the closure stresses p,(r) of all reinforcements with a weight function g(c, r) appropri-
-0 .J 7o
8'o
9'0
1do
14 o
12o
Energy Loss (eV) Fig. 6. Typical electron energy-loss spectra of the AI L edge. By subtracting the two reference spectra measured in the metal and in the ceramic from the interface spectrum, the interface component is obtained which contains information about the bonding between metal and ceramic.
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H. Prielipp rt al. / Materials
Science and En#wering
Al97
(1995) 19-30
Combining Eqs. (4) and (6), the fracture toughness at the plateau of the R-curve KS,, given at a crack length c,, where the first active bridge fails, can be written as KS = ]Ki + EL RI,. matrix(c) + EL R,,, mrtaJ”*
(7)
where K,, = (R, EL)‘!* is the crack tip toughness of the composite. SEM fracture surface observations reveal that in all composites with metal content greater than 20 vol.%,, alumina grains fracture essentially in a transgranular manner. Transgranular as well as intergranular fracture is observed only with a metal fraction of 10 vol.%. Accordingly, the toughening contribution by grain bridging can be neglected, and Eq. (7) can be reduced to K, = [K:, + KR,,.
Fig. 7. Representative fracture surfaces coarse-grained (b) AI!A120, composites 2, Al.
.* s
of medium-grained (a) and with 25 vol.‘%, Al: I, AIZOq;
R,,,(c) results from the J-integral formalism [36] (Eq. (5)) where u* is the opening at the last active bridge: R,,, = 2j;
I” P,(u)du 0
The stress displacement function p(u) uniquely describes the reinforcement characteristics [18] and is responsible for the increasing fracture resistance with increasing crack length R(c). The toughening contribution by plastic deformation of the ductile phases is governed by the mechanical properties of the metal, the interface properties (debonding), the metal ligament diameter and its volume fraction j In the case of intercrystalline fracture of the matrix, interlocking grains will remain in contact across the crack faces and possibly form bridges providing additional crack closure stresses [37,38]. In the elastic case or elasticcplastic case under smallscale yielding conditions, we can use the relation between applied stress intensity factor KA and mechanical energy release rate G, (with EL Young’s modulus under plane strain conditions):
,,,cta,(~)l’~~
(9)
Furthermore, TEM and EELS investigations predict strong AlPA120, interfaces independent of metal ligament diameter. In addition, the metal reinforcements exhibit scale invariant plastic deformation. Under the assumption that the mechanical properties are also independent of ligament diameter, the toughening contribution of the ductile phase is only governed by geometry, i.e. volume fraction and ligament diameter given metal/ceramic system.
800 600 400 200 0 Porosity 25 vol-%
Al 25 vol-%
Porosity 25 vol-%
Al 25 ~01-96
Fracture strength (a) and fracture toughness (b) of AI/AI,O, composites with 25 vol.% metal and varymg hgament diameter compared with the corresponding values for the porous preforms.
H. Prielipp et al..' Materials Science and Engbteerh~g, A 197 (1995) 19 30
25
compared with an alumina prepared from fine-grained A1203 powder of similar medium grain size. The fracture strength decreases from 560 to 100MPa with decreasing porosity (Fig. l l(a)) following an exponential relationship according to Rhyskewitsch [32]: c~t.= ao e x p ( - n P )
(1)
where n is an adjustable parameter, P is the fractional porosity and a0 is the fracture strength of the dense alumina. It is interesting to note that the fracture strength of A1 reinforced A12 03 is even higher than that of dense alumina. In the range 10 25 vol.% A1, the fracture strength reveals a constant value of about 700 MPa. With the metal fraction increasing from 25 to 40 vol.%, the strength increases approximately linearly to 800 MPa. The fracture toughness (Fig. l l(b)) decreases linearly with increasing porosity from 3.6 to 0.8 MPa m 1'2 for the alumina with 40% porosity. The fracture toughness increases monotonically with increasing metal content, up to 7 MPa m 1'2 at 40 vol.% A1. This is an improvement of nearly 100% compared with the dense alumina. c composites exhibit a similar trend, i.e. both the fracture strength and fracture toughness are enhanced (Fig. 12(a)) to 5 1 0 M P a (25vo1.% AI) and 650 MPa
Fig. 9. Optical micrographs of medium-grained AI/AI20 ~ composites with (a) 10 vol.% and (b) 40 vol.% AI.
3.3. Roh, o f metal content
Optical micrographs of alumina with medium ligament diameter infiltrated with 10 vol.% and 40 vol.% aluminum are shown in Fig. 9(a) and (b). The microstructures are homogeneous and it appears that a transition from islands of metal to islands of ceramic takes place (at least in a two-dimensional view). The average ligament diameter, as measured for polished sections, is also approximately a factor of 10 larger than the intrusion pore size (cf. Section 3.2). Fig. 10(a) and (b) compares fracture surfaces of the microstructures shown in Fig. 9. At low metal contents the ligaments deform by necking to a sharp point (separate ligaments), while at higher metal contents deformation to chisel edges (connected ligaments) takes place. For composites with metal content greater than 20 vol.%, fracture is essentially transgranular. With a metal content of 10 vol.% transgranular as well as intergranular fracture was observed (Fig. 10(a)). The results of fracture strength and fracture toughness of the medium-scale materials vs. porosity and volume fraction of A1 are shown in Fig. l l(a) and (b). In order to minimize the influence of grain size on fracture strength and toughness, the m composites are
Fig. 10. Fracture surfaces of medium-grained AI/AI203 composites with (a) 10 vol.% and (b) 40 vol.% AI.
26
H. Prielipp et al. /Materials Science and Engineering A 197 (1995) 19 30 1000
r
medium grained
r
0- porous
800 rl
:E
"r
800 ~.
[3 - i n f i R r a t e d
650-+ 6(
coarse grained ,
~ i ~,
600
t(-
400
.o,,o
....... "~. . . . . . ~ . . . . . . . . . .
200
-
O" = 560 exp(-4 P)
o
-'
/ dense
10 a)
20
40
30
a)
EL
medium grained [ O-
porous
o_
D - infiltrated t~
D-
[3
m
0
12 ¸ 10 / 8
AI content
10.5 +-0.7 Coarse grained 7.4±0.5
/
6
/
cm
2
/
~-
0
u~
~ i3!ii:ill
4.6t 0.5 i" /
4
iiiii!ii i;i~i~il;_ i
1.St 0.2
(,9 09
~-
Porosity
Porosity/AI content [vol-%]
8
~
7 25 vol-% 35 vol-%
25 vol-% 35 vol-%
............ "~"..... 5"...........
2
CO
i
0 o
b)
10
i
20
i
30
~:~:~:~:~:
::i:::i:::i:i dense
o
1.5t 0.1
25 vol-% 35 vol-%
25 vol-% 35 vol-%
i
40
Porosity/AI content [vol-%]
Fig. 11. Fracture strength (a) and fracture toughness (b) of AI/AI303 composites with medium ligament diameter and varying metal content compared with the corresponding values for the porous preforms.
(35volY0) and to 7 . 4 M P a m ~/2 (25vo1.% AI) and 10.5 MPa m 1/2 (35 vol.% ) (Fig. 12(b)). The strength and toughness of monolithic alumina with a grain size of 4 mm decreases with increasing porosity from 370 MPa for the dense alumina to 70 MPa for coarse-grained alumina with a porosity of 35% and from 4.6 to 1 . 5 M P a m ~/2 respectively. In comparison with the dense alumina, the A1 reinforcement yields an improvement in strength and toughness of more than 100%. Independent of metal ligament diameter, the Young modulus decreases from 260 GPa to 230 GPa for 25 and 35 vol.% AI respectively. Poisson's ratio remains roughly constant at 0.28. With E = 400 GPa for alumina, E = 70 GPa for A1 and a Poisson's ratio of 0.28 for both, the Young modulus can be evaluated for all the composites, using the arithmetic median of Paul's lower and upper Young's modulus of the composite. The computed data show a relatively good agreement with experimental measurements.
b)
Porosity
AI content
Fig. 12. Fracture strength (a) and fracture toughness (b) of A1/AI20~ composites with coarse ligament diameter and varying metal content compared with the corresponding values of the porous preforms.
as grains can reinforce a ceramic matrix by providing closure stresses in the crack wake shielding the crack tip from the applied stress. Process zone shielding [33] and crack deflection [34] are further possible mechanisms; however, they have only small toughening potential in these composites when compared with crack bridging [8,35].
4.2. Long crack toughness Observations on fracture surfaces of different scale microstructures point to scale invariant metal deforma-
O"
crack
4. Discussion
(I 4.1. Fracture mechanics Possible reinforcement mechanisms in metal ceramic composites are shown in Fig. 13. Ductile phases as well
~iiiY
Fig. 13. Schematic diagram showing possible crack bridging mechanisms in metal ceramic composites: bridging by ductile phases (1) and bridging by matrix grains (2).
H. Prielipp et al.
Materials Science and Engineering, A 197 (1995) 19 30
_rntai,,x[ crac~
coarse
medium
a)
C
R
P
.**/o.-- ........ ,."
m
f U
b)
c)
Fig. 14. Schematic diagram describing p(u) functions (b) and Rcurves (c) for metal reinforced ceramics with varying ligament diameter (a), based on the assumption of scale independent metal deformation.
27
metal is a linear function of volume fraction 1/1 as well as of metal ligament diameter d owing to the microstructural scale invariance of plastic deformation. This implies according to Eq. (9) that variations in fracture toughness with metal content should obey a square root dependence on Vf for a given microstructural scale of the metal phase. Fig. 15 demonstrates good agreement between theory and experiment, except for the coarse-grained composite with a metal content of 25 vol.%. In order to estimate R~...... ,,~ from Eq. (5), the p(u) function was assumed to be constant up to a m a x i m u m crack opening where the first ligament fails. According to literature results [10] this opening is about the same size as the ligament diameter, which in turn can be estimated from optical micrographs to be roughly a factor of 5 larger than the pore channel size measured with mercury porosimetry. Thus the solid lines of Fig. 15 correspond to ligament diameters of 0.4, 1.25 and 4 jam. Combining Eqs. (9) and (5) results in a closure stress for the AI of 400 MPa, which exceeds the yield stress of pure AI by an order of magnitude. The authors relate this fact to the constrained plastic deformation as well as the minute microstructural scale of the single-crystal ligaments.
4.3. Composite .fracture strength tion and debond lengths. Consequently, p(u) functions can be derived from one assumed function and transformed into another microstructural scale simply by stretching the crack opening parameter in accordance with the ligament diameter (Fig. 14). Large reinforcements will therefore lead to smaller crack closure stresses at given small COD, which will be maintained up to higher crack opening displacements. This will result in R-curves for large reinforcements which can be characterized by small initial slope with a high peak thoughness. Therefore, the long-crack fracture toughness (as measured by SEPB) will increase with ligament diameter. The SEPB method yields fracture toughness values for bridged cracks of lengths between 200 and 400 jam, which for our materials is near the plateau of the R-curve and can therefore be represented by Eq. (9). The crack tip toughness Ko can be determined by a procedure [18] which relies on a measurement of near crack tip opening displacements and use of the crack profile description of Barenblatt [39]. A crack tip toughness of 2 . 0 M P a m 12 for A1/A1203 composites produced by direct metal oxidation has been estimated [40] which can be compared with values of 2.0 MPa m ','2 for alumina [18]. Fig. 15 shows the results of the fitting procedure according to Eq. (9), taking Ko as 2.0 MPa m ~'2 for all composites and assuming that crack bridging is due to ductile reinforcements only (Section 4.2). The sole assumption required is that the fracture energy contribution R t..... lal of the ductile
The increase in fracture strength with metal content may seem surprising at first, since AI has a yield strength much below the strength of alumina, thus any rule of mixture does not apply. To understand the strength of metal/ceramic composites, one needs to appreciate that, in interconnecting networks, the ceramic provides the failure site at the largest flaw, and the metal provides the fracture toughness. In the following sections, qualitative descriptions are provided for trends of composite strength, with the assumption that the initial flaw size of the porous alumina is the same as that of the respective composite.
12 &- small I • - medium I m
8
¢/) ¢/)
6
0-
o~ c¢,=
4
c~ "~ o I-
2
~
-e ~ EqarSe
± ~
_
m
0 0
10
20
30
40
AI content [vol-%] Fig. 15. Influence of metal ligament diameter and volume fraction on steady state toughness of AI/AI20~ composites. Full lines are representations of Eq. (9).
28
H. Prielipp et al. / Materials Science and Engineering A 197 (1995) 19 30
Gm
Gs Gc
..... // //
G,R //
////
y('*
~ " "
............ /,.::~<:'.:~........................... a c
...../~, "/ ..,, . / ~ "
//
increased metal fraction
//].~#~"
.,"/ ,~! • ........~>..,' /
G, R Rm
/-~ /
.....' .";'::!';;;" ......., " "~.~:.......................................
/ " ,.,;.S,I ,,**
Rs
....::..::-IS .....
IL
C
C Fig. 16. Schematic diagram describing R-curves of metal reinforced ceramics with varying metal ligament diameter. The tangency condition of the applied energy release rate GA (straight, dashed lines) with the R-curves yields the respective instability points.
4.3.1. Effect of metal ligament diameter In describing the fracture strength, the actual slope of the R-curve, as well as the initial flaw size, are important factors. Inserting the values for toughness, as measured with SEPB, and strength in Griffith's equation and solving for the crack length yields almost identical flaw sizes for the s, m and c alumina with 25% porosity. In the schematic diagram in Fig. 16, identical initial flaw sizes are therefore used for all composites. The c composite exhibits a rather shallow R-curve so that crack instability occurs with a large degree of stable crack growth; hence in spite of a large plateau toughness the strength is rather low. The finegrained micostructure exhibits only a very small Rcurve height, so that the increase in strength is not as dramatic. Conversely, the m material appears to combine initial flaw size and slope of the R-curve most favorably and therefore has the highest fracture strength of 710 MPa. Due attention must be paid to the scale of the initial flaw size. In cases where this parameter is increased (as in thermal shock), the large-scale microstructure may prove to be advantageous compared with the fine- or medium-scale counterparts [41].
Fig. 17. Schematic diagram describing R-curves of metal reinforced ceramics with varying metal content. The tangency condition of the applied energy release rate GA (straight, dashed lines) with the R-curves yielded the respective instability points.
4.3.3. Effect of flaw bridging Our qualitative considerations describe only the trend of strength as a function of ligament diameter as well as of volume fraction. Assuming the same initial flaw sizes for the composites as for the corresponding porous alumina, the instability criterion (tangency condition of the applied stres intensity factor with Rcurve) predicts an enhancement of the composite strength directly proportional to the fracture toughness at the instability crack length. This statement is consistent with the initial modeling from Bao and Zok [21]. Quite on the contrary, however, our results show a larger enhancement of strength than of toughness owing to metal infiltration, as demonstrated in Fig. 18 by the mechanical property data normalized to the data of the porous preform. The above model assumes that the initial flaw is not bridged by the ductile phase. Based on investigations on the fracture strength of alumina [24], where the largest pore determines the failure origin, we assume that the site of the initial flaw remains unchanged but is converted to a metal filled cavity. Therefore the crack is initially bridged by metal with a volume fraction fb of 10
4.3.2. Effect of metal content A large metal content results from a preform with high porosity. In applying the Griffith equation, it appears that the initial flaw size increases with increasing porosity. Depending on the exact interaction of initial pore size and slope of the R-curve, an increase or decrease in strength with metal content may result. In the case of AI reinforcement, the Rcurves are steep enough to compensate for the small increase in initial flaw size and therefore yield an increase in strength with increasing metal content as shown schematically by the tangency condition in Fig. 17.
medium grained © - Strength [] - Toughness
~8
.o .. O' /' . .O" ....O ......
~(
~
2
,:::: ::::::?
El " "
. ..,,D ..
....
......
10
20
30
40
AI content [vol-%]
Fig. 18. Effect of AI content on the relative enhancement of the composite strength and toughness. The composite properties were normalized to the corresponding values for the porous alumina preforms.
29
H. Prielipp et al./ Materials Sciem'e and EnghwerhTg A 197 (1995) 19 30
long crack
metal
finitial : fb = 1 0 0 %
,/'
G, R
// a)
1
.
G,R
/,/'/ / c
b)
c
Fig. 20. R-curves (a,b) explain schematically the short crack anomaly (b) compared with long crack behavior (a).
fb = F(C) > f
1
small crack
fb = f = const.
C
Fig. 19. Hypothesis for crack initiation in metal/ceramic composites where failure-causing pores are eliminated and failure initiates at large metal filled cavities.
100% (Fig. 19(a)). As the crack extends into the matrix, it will consecutively encounter an increasing number of metal ligaments (Fig. 19(b)). If the crack length is large with respect to the metal filled cavity, the metal bridging fraction is constant (f, = J ) (Fig. 19(c)). Inclusion of the change in metal volume fraction leads to a modification of Eq. (5), where now f has to be replaced by
tion into tailored porous preforms. These materials exhibit fracture toughnesses up to 10.5 MPa m j'2 and fracture strength up to 810 MPa. In contrast to metal matrix composites (e.g. [42]), both the fracture strength and fracture toughness are increased with second phase content. TEM and EELS characterization show continuous, microcrack-free bonds in all composites of varying microstructural scale. Long-crack fracture toughness data can be inferred by using the plateau toughness of one sample microstructure and then scaling the toughness with respect to ligament diameter and metal volume, where the observation of scale invariant plastic deformation is utilized. Trends in fracture strength as a function of metal content and metal ligament diameter can be explained. A striking result, however, is the fact that the fracture strength increases more pronouncedly by metal infiltration than the plateau toughness derived from long crack measurements. This fact is attributed to large closure stresses during and after crack initiation at metal filled cavities associated with plastic deformation of these large metal spheres.
f(c): t¢i*
R,,= 2f(c)
Ij
pi(u)du
(10)
The foregoing discussions on the influence of the matrix on fracture are therefore still valid, but have to be augmented by the closure effect of the large metal filled pore. In addition, the initial slope of the R-curve is steeper than that of a long-crack R-curve (Fig. 20(a) and (b)). This qualitative description, by separating small- and long-crack fracture toughness, can explain the fact that the composite strength increases more than the (measured) long crack fracture toughness would predict (tangency condition in Fig. 20(b)).
Acknowledgements
We thank Maria Sycha for preparing the TEM specimens and Christina Scheu for performing part of the experiments and analysis. S.K.S. wishes to thank the Alexander von Humboldt-Stiftung for financial support. This work was supported by the Deutsche Forschungsgemeinschaft under contract number Ro 954/1-1 and the Volkswagen Foundation under contract number 1/66 760.
References 5. Conclusion
A1/AI203 composites can be produced with widely differing microstructures (especially metal volume and metal ligament diameter) by gas pressure metal infiltra-
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H. Prielipp et al. / Materials Science and Engineering A 197 (1995) 19-30
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[28] H. Mfillejans and J. Bruley, Ultramicroscopy, in press. [29] M.C. Bhardwaj, Am. Ceram. Soc. Bull., 69 (9) (1990) 1490. [30] F. Zok and C.L. Holm, Acta Metall. Mater., 38 (10) (1990) 1895 - 1904. [31] M. Knechtel, H. Prielipp, N. Claussen, S.K. Streiffer, C. Scheu, M. Rfihle and J. R6del, to be submitted to J. Am. Ceram. Soc. [32] Rhyskewitsch, J. Am. Ceram. Soc., 36 (2) (1953) 65-68. [33] A.G. Evans and R.M. Cannon, Acta Metall., 34 (5) (1986) 761 800. [34] K.T. Faber and A.G. Evans, Acta Metall., 31 (1983) 577-584. [35] J. R6del, J, Eur. Ceram. Soe., 10 (3) (1992) 143 150. [36] J.R. Rice, J. Appl. Mech., 35 (1968) 379-386. [37] R. Knehans and R.W. Steinbrech, J. Mater. Sci. Lett., 1 (1982) 327 329. [38] J. R6del, J. Eur. Ceram. Soc., 9 (4) (1992) 323 334. [39] G.I. Barenblatt, Adv. Appl. Mech., 7 (1962) 55-129. [40] J. R6del, M. Sindel, M. Dransmann, R.W. Steinbrech and N. Claussen, J. Eur. Ceram. Soc., 14 (1994) 153-161. [41] S. Sch6n, H. Prielipp, R. Janssen, J. R6del and N. Claussen, J. Am. Ceram. Soc., 77 (3) (1994) 701 704. [42] I.A. lbrahim, F.A. Mohamed and E.J. Laverna, J. Mater. Sci., 26(1991) 1137-1156.