SiCw composite

ItATERIALS SCIEMCE& E_IMEERIMe ELSEVIER

A

Materials Science and Engineering A209 (1996) 251 259

A microstructural investigation of the mechanisms of tensile creep deformation in an AI 2 0 3 /SiC w composite C. O'Mearaa, T. Suihkonena, T. Hansson b , R. Warren

C

"Department of Physics, Chalmers Unirersitr of Technology, Giitehorg S-412 96, Sweden bDepartment of Mechanical Engineering, Nagaoka University of Technology, Nagaoka, Japan "Department of Materials Science and Production Technology, Luled Unitwsitr of Technology, Luled, Sweden

Abstract The tensile creep behaviour of an SiC w (25';;;1) reinforced alumina composite was investigated using scanning electron microscopy (SEM) and transmission electron microscopy and automatic image analysis in SEM. The creep tests were carried out in air in the ranges 1100-1300 °C and 11-67 MPa. Each creep test was performed at a constant temperature. The material had a stress exponent of about three for all temperatures and an approximate activation energy of 650 kJ mol I. The creep resistance of this composite is poorer than that of similar composites studied earlier. Microstructural examination revealed the microstructure to be extremely inhomogeneous consisting of spherical whisker-rich clusters (20-100 ,urn) surrounded/separated by AI 20 3 rich rims (10 ,urn). The secondary creep rate is dominated by a damage accumulation process namely cavitation and crack growth in both the SiC clusters and the AI 2 0 3 rims. Final fracture seems to occur through the alumina rich regions. The lower creep resistance of this composite compared to that of similar composites is attributed primarily to the inhomogeneity of the as-received material. Keywords: Tensile creep deformation; Alumina composites; Microstructure

1. Introduction Monolithic alumina exhibits only moderate strength and creep resistance and, like most monolithic ceramics, is extremely brittle. SiC whisker reinforcement of alumina (SiC w /AI 2 0 3 ) has been employed primarily and successfully to improve fracture toughness [1-4]. Various toughening mechanisms such as whisker bridging and pullout, microcracking and crack deflection are operative depending on microstructural factors and experimental conditions [1-4]. The improvements obtained in the composite in both fracture toughness and strength as compared to monolithic alumina has led to its application as, for example, cutting tool inserts and extrusion valves, and give it potential for use in structural applications at high temperatures. However the use of ceramic materials in high temperature structural applications is inevitably dependent on their time dependent mechanical properties such as creep and oxidation resistance. The incorporation of SiC0921-5093196/$15.00 © 1996 SSDI0921-5093(95)10102-0

Elsevier Science S.A. All rights reserved

whiskers into alumina is also expected to improve the creep resistance and this has largely been confirmed in bend and compression tests [5-12] but not in tension [13]. Observed creep mechanisms in monolithic alumina include basal slip, diffusional creep and grain boundary sliding (GBS) [14,15]. Grain boundary cavitation has also been observed in association with GBS [16,17]. The stress exponent has generally been found to vary between I and 2 and the activation energy between 400650 kJ mol- I. The creep resistance has been found to increase with grain size and there is a general agreement that aluminas with "clean" grain boundaries exhibit higher creep resistance than aluminas with an amorphous grain boundary phase [14-24]. Reinforcement of alumina with SiC-whiskers is expected to improve the creep resistance primarily by interlocking/pinning of grains which then limits grain boundary sliding [8]. However, several factors can be identified which will affect the mechanical response of

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the composite under creep conditions: the volume fraction of whiskers; the strength of the interfacial bond between the fibre and the matrix; the grain size of the matrix grains; the volume of intergranular amorphous phase and; the oxidation susceptibilty of the material which causes the formation of glass at the whisker matrix interface. These factors will vary from material to material and will complicate both the interpretation of the creep behaviour and the comparison of different works. Studies on bending and compression creep of the composite indicate that there exists a transition stress below which the creep is dominated by diffusion accommodated mechanisms and above which a damage accumulation process involving cavitation and microcracking become increasingly important. In the low stress regime the stress exponent is 1-2 while above the transition it increases to values of 5 [6-10]. Activation energies similar to those in monolithic alumina are observed. No direct evidence of dislocation activity has been found. In general whisker reinforcement increases the creep resistance of alumina, however in the high stress regime whisker volume fractions above 20% do not provide further improvement and may even decrease the creep resistance [8]. Grain boundary and interfacial amorphous phases are detrimental to creep resistance and may promote transition to a damage accumulation process [9]. This is consistent with the observation that creep resistance is lower in air than in inert atmospheres since the composite is sensitive to oxidation. Work by the authors on tensile creep of the composite has shown that in tension even in the low stress regime damage accumulation was the dominant creep mechanism and a stress exponent of three was obtained for all temperatures and stresses [13]. Previous investigations on ceramic materials tested in tension and flexure have shown stress exponents of three to arise from creep cavitation [25]. However for composite and multiphase ceramics because of the complex interaction between the microstructural constituents, creep deformation mechanisms cannot be reliably deduced from creep data alone [9] but require direct observation of the deformed microstructures. This work presents a microstructural examination of the tensile creep behaviour of a SiC w (25%) reinforced alumina composite. Electron microscopy was used to obtain information on the possible creep mechanisms.

The mixture was cold pressed to 20 mm diameter rods and then HIPped (1600 DC, 160 MPa, 1 h). Cylindrical test specimens were produced by precision machining from the rods. Each specimen had a total length of 150 mm and a diameter of 10 mm reducing to 4 mm over a 20 mm long gauge length. 2.2. Creep testing The creep equipment was specially designed for the testing of brittle materials and details of the test system are described in Ref. [13]. The creep tests were carried out in air in the ranges 1100-1300 °C and 11-67 MPa. Each creep test was performed at a constant stress and temperature. Two specimens were pre-heat treated in air at 1300 °C prior to creep testing. The high temperature heat treatment was used to investigate the effect of oxidation on the creep behaviour by comparing with non heat treated specimens subjected to the same creep conditions. 2.3. Microstructural examination The microstructure of the as-received and crept materials were studied using both scanning and transmission electron microscopy (SEMjTEM) and quantitative SEM using automatic image analysis (AlA). 2.3.1. TEM Thin sections for TEM analysis were taken from the centre of the gauge section directly above the fracture surface and were cut in the longitudinal direction, parallel to the stress axis. Thinned sections were dimple ground followed by ion-beam thinning to perforation. TEM examination was carried out using a JEOL 2000FX TEMjSTEM instrument equipped with a Link Systems AN 10000 EDX spectrometer. 2.3.2. SEM SEM specimens were cut from the gauge section in the longitudinal direction from the fracture surface to a distance of about 0.5 em along the gauge length. The specimens were mounted in transoptic plastic, ground and polished down to 0.25 jim using a Struers semiautomatic polishing apparatus. The specimens were examined in a CAM Scan S-4 80DV instrument equipped with a Link Systems AN 10 000 EDX spectrometer. The specimens were examined in secondary and backscattered electron mode.

2. Experimental 2.1. Material The composite was produced from a powder mixture of alumina and whiskers without sintering additives.

2.3.3. Quantitative microscopy (AlA) was used for cavity, alumina grain size and phase volume fraction estimation. A Jeol JXAj8600 Electron Probe Micro Analyser (EPMA) was used with Kantron software.

C. O'Meara et al. Materials Science and Engineering A209 (1996) 251-259

253

Table 1 Tensile creep test conditions and results Stress (MPa)

Time to fracture (h)

(';i,,)

2"

1200 1100

19 11-35

42.1 728

3.4 2.3

3 4 5 6b 7 8 9 IOc

1200 1200 1300 1300 1300 1300 1275 1300

35 67 19 11 19 35 35 11

13.9

1.7

1.2

1.2

I

Temp

Strain to fracture

(0C)

Sample

1.7 252 13.7 0.17 2.2 4.9

2.5 3.4 3.0 1.5 2.5 1.3

Secondary creep rate (s 1)

Creep exponent

1.9 x 10- 7 3.7 x 10- 1°_ 1.6 x 10 R 2.5 x 10- 7 2.4 X 10- 6 3.8 x 10 6 2.0 x 10 " 4.1xlO- 7 2.3 x 10- 5 2.8 x 10 6 6.6 x 10- 7

3.25 3.25

Preheated at 1300 °C (h)

3.28 3.28 2.94 3.7 3.08 2.94 3.08 3.08

61 72 72

" Load increased during the test. b Test aborted before failure. c Defect in the sample.

2.3.3.1. Cavity analysis. Two sets of measurements were made for each fractured specimen, one at the fracture surface and one 3 mm further into the bulk. At a magnification of 5000 x, 50 fields were examined in each measurement with a total frame area of 10 x 5 x 327.12 j1m 2 = 16356 j1m 2 • Cavities/pores in the size range 0.1 -10 j1 m were measured by AlA. The parameter measured was the area of the pore. This was converted to the equivalent diameter of that area by DCIRCLE

=

2

ft

(I)

Where DCIRCLE is the average diameter of the pore and A is its measured area.

2.3.3.2. Grain size analysis. For the alumina grain size analysis the specimens were first etched in argon at 1300 °C for 15 min. For each specimen 10 micrographs at a time were taken in a line close to and parallel with the fracture surface. At least 500 alumina grains were analysed in each specimen. The parameter measured in AlA was the area of the grain which was converted to average diameter by Eq. (1). 2.3.3.3. Volume fraction of phases. As will be discussed in the results, considerable inhomogeneity in the matrix was observed dividing the microstructure into "whisker/rich" and "alumina rich" areas. A qualitative AlA analysis was carried out simply by marking the "whisker/rich" areas in one colour and the "alumina rich" areas in another in order to get an indication of the extent of inhomogeneity. In addition an estimation of the whisker fraction in the "whisker/rich" and "alumina rich" areas was carried out using EDX analysis.

3. Results

3.1. Creep results A summary of the creep results is given in Table 1. The composite exhibited limited regions of primary creep, relatively well defined secondary creep and entered a tertiary creep region just before fracture for all conditions of stress and temperature. The material had a stress exponent of approximately three for all temperatures and an approximate activation energy of 650 kJ mol- '. The creep resistance of this composite in tension is poorer than that of similar composites studied earlier in bend or compression. However, the creep resistance improved significantly following high temperature pre-heat treatment.

Fig. 1. SEM image showing the inhomogeneity in the microstructure, spherical whisker-rich clusters surrounded/separated by AI 20, rich rims.

c.

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O'Meara et al. / Materials Science and Engineering A209 (1996) 251-259

Table 2 The cavity size measurements of cavities between 0.1 and 10 .urn Sample and location

Area density of cavities (\/mm 2 )

Median cavity diameter (.urn)

Mean cavity diameter (.urn)

Cavity nucleation rate (\ jh)

Area fraction of cavities CY<,)

0 1 2 3 4 5 6 7 8 9

22 liS 170 178 96 200 163 126 198 161

0.42 0.50 0.52 0.40 0.30 0.51 0.50 0.51 0.53 0.41

0.50 0.60 0.60 0.50 0.38 0.60 0.56 0.61 0.57 0.45

45 4.0 210 1306 1980 10.57 ISO 19317 1214.9

0.42 3.26 5.0 3.10 l.l 5.6 4.0 4.0 5.1 2.6

o stands for as-received

sample.

3.2. Microstructure 3.2.1. As-received material As is shown in Fig. 1 microstructural examination revealed that the as-received material had an inhomogeneous microstructure consisting of spherical whiskerrich clusters (~20-100 ,urn) surrounded/separated by AI 2 0 3 rich rims (~ 10 ,urn). The inhomogeneity is probably inherited from spray drying used during processing of the composite powders which resulted in a clustering of the SiC phase. The results of AlA and EDX analysis indicate that approximately 40% of the microstructure has a SiC content of less than 15 vol.% while about 60% has a SiC content of '" 30 vol.%. The microstructure is in effect a composite within a composite. In addition the alumina grain size was small (sub-micron) and bi-modal in distribution with the larger grains in the rims and very small grains in the whisker clusters. In TEM very little ( < I nm) intergranular amorphous phase could be detected in the microstructure. The alumina grain morphology was variable with no evidence of preferred orientation. Very few dislocations were observed but small pores within the AI 20 3 grains were detected. Whisker lengths of up to 10 ,urn were observed. No coating could be detected on the SiC whiskers but individual whiskers were uneven in diameter with micro-roughness resulting from the stacking of SiC polytypes. A small amount of porosity, up to 0.5 ,urn in diamter at Al 20 3 triple grain junctions and at AI 2 0 3 /SiC and SiC/SiC interfaces was observed and corresponds with the 0.4% porosity estimated in SEM by AlA (see Table 2). 3.2.2. Crept material SEM and TEM examination of the crept materials revealed that all samples contained two distinct families of creep damage (i) a large volume of sub-micron and micron size (0.1-10 ,urn) pores and cavities and (ii) cracks of;:::: 10 ,urn in length. (i) The cavities/pores were observed at AI 2 0 3 /AI 2 0 3 ,

AI 2 0 3 /SiC and SiC/SiC contacts. The extent of cavitation decreased somewhat with distance from the fracture surface but was evident the entire length of the specimens (~5 ,urn). Cavitation between the Al 2 0 3 grains was observed to be very frequent and to occur both at triple and two grain junctions. Cavitation occurred most commonly at two grain contacts and here the grain facets were separated the entire length of the grains forming lath like cavities (Fig. 2). The nucleation of smaller cavities along two-grain facets was also frequently observed in TEM. These cavities were approximately 100 nm in length and lenticular in shape and encroached on both grains as is shown in Fig. 3. No grain growth was detected following creep even for the preheated samples and few dislocations were observed. Indication of possible GBS was found both in TEM and in SEM examination of etched specimens (see Fig. 4). Cavities also formed at AI 2 0 3 /SiC and SiC/SiC contacts. As is shown in Fig. 5, SiC debonding and pullout occurred frequently. When these debonded contacts lay close together a small isolated intergranular cavity formed. Fracture of long SiC whiskers was also observed. Fig. 6 shows the presence of an amor-

Fig. 2. TEM bright field image showing a lath like cavity between two AI 2 0} grains following creep testing at 35 MPa and 1200 0c.

C. O'Meara et al.

255

Materials Science and Engineering A209 (1996) 251-259

"t

'~ ... :j.. "

',ii "

Fig. 3. Cavities along two-grain A1 20, facets in a specimen creep tested at 1300 DC and 35 MPa (TEM).

phous coating (~2 nm) observed in TEM on the surface of some SiC whiskers in cavitated areas. This phenomena was observed in all specimens examined indicating that these interfaces were oxidised during creep exposure, It is pointed out that the TEM foils were taken from the centre of the specimens where it is assumed that the effects of oxidation are at a minimum; nonetheless, some qualification must be made on these observations: (1) the amorphous film was not observed at all AI 2 0 3 /SiC and SiC/SiC cavitated contacts in the same specimen; (2) no significant increase could be detected in the volume of intergranular amorphous phase in non-cavitated regions of the specimen or at the A1 20 3 /A1 2 0 3 cavities as compared to as-received material; (3) Fig, 6, a typical oxidised cavity shows a cleanly fractured SiC whisker surrounded by an even layer of amorphous phase, no cavitation or glassy ligaments are observed in the amorphous phase which appears to "close off" the cavities. The authors therefore feel that oxidation of the whiskers may equally well have occurred after cavitation as much as being the facilitator of cavitation as suggested in other works. Quantitative analysis was undertaken to investigate differences in cavitation between the samples. The results for the area

/

Fig. 4. Grain boundary sliding of alumina grains observed in SEM.

r. ~.~

"'~I"~;"".", ..::cc.

_~""."i/

;!~"il

Fig. 5. TEM images showing whisker debonding (a) and pullout (b) in crept material.

close to the fracture surfaces are given in Table 2. All the specimens showed a similar pore/cavity size distribution with a mean cavity size of approx. 0.5 ,um. In all cases 9Y;;, of the cavities were under 1.5 ,um in length which agrees well with TEM observations. Even with the limited number of test points a linear-stress relationship at temperature is observed between both the area density of cavities and the cavity nucleation rate (total number of cavities/time to fracture) (Fig. 7). The pre-heat treated specimens show the same dependence but at lower values indicating some positive effect of the heat treatment. These results do imply that cavitation is the main deformation mechanism occurring during creep and that it was operative under all test conditions. However the results also imply that creep rupture does not occur when a similar level of cavitation has been attained but rather that the level of tolerable cavitation is highly stress/temperature dependent due to an interplay of additional microstructural factors as will be discussed later. As is shown in Fig. 8, more severe creep damage in the form of cracks ( ~ 10 ,um) was observed: (i) through the Al 2 0 3 in the rims around the clusters originating at the edge of the specimen and; (ii) in the interior of the clusters associated predominantly with the SiC whiskers. In both cases the cracking is mainly perpendicular to the tensile direction and is intergranular. The

256

C. O'Meara et al.! Materials Science and Engineering A209 (1996) 251-259

frequency of the cracks decreases with distance from the fracture surface. The distribution, type, size and severity of the cracks varied with temperature and stress. The cracks in the SiC clusters (Fig. 9) were found in all specimens but were most frequent in specimens tested at 1200 0c. These areas were also observed in TEM as interconnecting cavitated networks of SiC grains where cohesion is maintained by an amorphous coating on the surface of the SiC whiskers. In general these cracks remained isolated within the clusters, i.e. did not join up via the rims with neighbouring cracks. Thus these cracks were generally 20-50 ,um in length. As is shown in Fig. 8, the cracks in the Al 2 0 3 rich

10

~

b.

~

....=

,~

eu

0.1

Qj

1

Cj

= =

.....

....

0.01

'5=

eu

b.

•

•

0

0

--=

0

1200°C

b.

1300°C

••

b.

U

0

0.001 10

1300°C, preheated 61 h 1300°C, preheated 72h

I

I

I

I

I

20

30

40

50

60

70

Stress (MPa) Fig. 7. Cavity nucleation rate vs. stress.

regions originated predominantly at the edge of the specimens and were found to extend from 20 ,um up to half way through the specimen diameter and were responsible for creep rupture. As is seen in Fig. 10 the cracks also penetrated through these regions intergranularly. At whiskers lying in their path they were either deflected along the interface or remained bridged by the whiskers. Cracking through the Al 2 0 3 was most severe in specimens tested at 1300 0c. In the pre-heat treated specimens cracking was predominantly of this type with much fewer SiC cluster cracks developing. In specimens tested at 1200 °C a considerable amount of cracking or crack branching parallel to the tensile direction is observed. Examination of etched specimens indicated this cracking was frequently intragranular and was most severe in the specimen tested at 67 MPa. The preheated specimens and the specimen tested at the highest stress showed the most severe crack damage. The oxide scale thicknesses and extent of oxidation of the cracks were examined in SEM and although in general the scale thickness varied with length of exposure the scales are thicker at 1300 °C as compared to 1200 0c. However there was very little oxidation in the cracks from the surface in the pre-heat treated specimens again indicating some effect of pre-heat treatment on the microstructure.

4. Discussion

Fig. 6. TEM images showing the presence of an amorphous coating ( ~ 2 nm) on the surface of SiC whiskers in cavitated areas: (a) bright field and; (b) dark field (35 MPa, 1300 Qq.

The results of this work indicate that in tension the secondary creep is dominated by a damage accumulation process namely cavity and crack growth. This is consistent with a calculated stress exponent of approximately three for all test conditions and that of previous investigations on ceramic materials tested in tension and flexure where stress exponents of three have been found to arise from creep cavitation [25]. However the creep resistance is inferior to both monolithic alumina

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C. O'Meara et al. Materials Science and Engineering A209 (1996) 251-259

and to similar composites tested previously in bending or compression. In addition creep controlled by damage accumulation processes is occurring in this composite at stresses far below the threshold stress observed for this mechanism in compression and bending (60-170 MPa). Although the mode of testing, tensile as opposed to bending or compression, is expected to have some effect on the creep behaviour in particular to promote the damage accumulation process and reduce creep failure strains it is thought that this microstructural examination has provided microstructural explanations for this poor creep behaviour. The inhomogeneity of the microstructure means that the material is literally a composite within a composite

.

...1-..

•

"""!

.~

Fig. 9. Crack development in the whisker clusters is associated predominantly with the SiC whiskers. (11 MPa, 1100 ec, SEM).

and consists to approximately 60°;() of clusters with a SiC whisker content 30 volume which is well in excess of the optimum for creep inhibition of alumina. These clusters are surrounded by an interconnecting alumina rich "skeleton" which makes up approximately 40% of the microstructure and has a SiC content of less than 15 volume. Although this is a low fibre content other works have shown that additions of even 5 wt.'% SiC whiskers causes improvement in creep resistance [7]. In addition the alumina grain size is very small (largely sub-micron) which pre-empts a lower creep resistance and is also bi-modal in distribution with the larger grains in the rims and very small grains in the whisker clusters. Improved creep resistance in this composite is not necessarily to be expected with this micostructural scenario and the results obtained need to be examined in this context. Microstructural observations and quantitative analysis indicated that the extent of cavitation at failure has a definite stress dependence at temperature, the extent of crack growth and the volume of the two different types of crack was also stress dependent at temperature so that the level and type of damage accumulation at failure was different for the different specimens. The net

Fig. 8. SEM backscattered electron image showing severe cracking observed in crept specimens (a) through the AI 20, in the rims around the clusters originating at the edge of the specimen (19 MPa, 1300 ec) and; (b) in the interior of the clusters associated predominantly with the SiC whiskers. (19 MPa, 1200 ec).

Fig. 10. SEM image showing an intergranular crack in a rim following creep testing at 19 M Pa and 1300 ec.

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C. O'Meara et al. / Materials Science and Engineering A209 (1996) 251-259

damage accumulation will be a summation of damage effects in the two different elements of the microstructure. The clusters have a high SiC content and can be expected to cavitate easily as observed. The skeleton of Al 2 0 3 with low whisker should be more creep resistant and will act as the reinforcing phase encapsulating the SiC damage and preventing it from linking up. Cavitation and crack formation in the clusters would then be a major contributor to the secondary creep rate and SiC cluster cracks are observed to be more frequent in specimens with a high secondary creep rate. Cavitation also occurs in the Al 20 3 rich rims phase but not as easily as in the SiC clusters. Final fracture seems to originate from the surface in the alumina regions and then to propagate through these regions intergranularly either when the fracture stress is reached for alumina of this particular grain size (specimen 4) or when a sufficient number of contiguous cavities have been built up such that the crack is subject to a stress intensity K ~ K rh [26] (specimens 1~3, 5-9). This interpretation is consistent with and can explain the differences in crept microstructures, times to failure and secondary creep rates of the tested samples. The effect of oxidation on the creep behaviour was difficult to evaluate as all the samples had been subjected to high temperature air exposure. Although other microstructural factors are responsible for creep failure, with reference to other work it is highly likely that oxidation during creep exposure does facilitate cavitation. In this work some oxidation of the SiC whiskers in the matrix was observed in TEM although no significant increase could be detected in the volume of intergranular amorphous phase in non-cavitated regions of the microstructure as compared to as-received material. The TEM evidence suggests that oxidation of the whiskers had occurred after cavitation and may bond the whiskers or the whiskers and the matrix grains together. In addition the specimens which were pre-heat treated had a lower density of cavities, lower cavity nucleation rates, fewer SiC cluster cavities and lower secondary creep rates than specimens subjected to the same conditions without pre-heat treatment. In general the specimens with the lowest secondary creep rates were also those which were the most severely oxidised. At this stage it is therefore difficult to accurately identify the role of oxidation, whether positive or negative, on the creep deformation process(es).

5. Conclusions The tensile creep behaviour of a SiC (25%) reinforced alumina composite was investigated in air in the ranges 1100~ 1300 °C and 11-67 MPa. The as-fabricated microstructure was found to be extremely inhomogeneous consisting of spherical whisker-rich clusters (20-

100 ,um) surrounded/separated by Al 2 0 3 rich rims (10 ,um). Stress exponents of approximately three correspond well with the microstructural analysis which indicated that the secondary creep rate is dominated by a damage accumulation process namely cavitation and crack growth in both the SiC clusters and the AlP3 rims. Final fracture seems to occur through the alumina rich regions. The poorer creep resistance of this composite compared with that of similar composites is attributed primarily to the inhomogeneity of the as-received material.

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of fine-grained alumina (AllO,): 11, Basal slip and nonaccomodated grain-boundary sliding, J. Am. Ceram. Soc., 63 (1980) 53. [23] J.R. Porter. W. Blumenthal and A.G. Evans, Creep fracture in ceramic polycrystals-L Creep cavitation effects in polycrystalline alumina, Acro Merall" 29 (1981) 1899. [24] R.A. Page. J. Lankford and S. Spooner, Small-angle neutron scattering study of creep cavity nucleation and growth in sintered alumina. J. MareI'. Sci., /9 (1984) 3360. [25] D.F. Carroll and R.E. Tressler, Accumulation of creep damage in a siliconized silicon carbide, 1. Am. Ceram. Soc" 7/ (! 988)

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