Creep behaviour of SiCAl materials

Materials Science and Engineering, 71 (1985) 147-157

147

Creep Behaviour of SiC-AI Materials* J. L. CHERMANT and F. OSTERSTOCK Equipe Mat~riaux-Microstructure du Laboratoire de Cristallographie, Chimie et Physique des Solides, Laboratoire associ~ au CNRS 251, Institut des Sciences de la Mati~re et du Rayonnement, UniversitY, 14032 Caen C~dex (France)

(Received October 22, 1984)

ABSTRACT Various hot-pressed SiC alloys w i t h alum i n i u m have been investigated using creep in t h r e e - p o i n t bending in the 1 6 7 3 - 1 9 7 3 K temp e r a t u r e range. Creep characteristics were analysed as a f u n c t i o n o f m o r p h o l o g i c a l p a r a m e t e r s and o f the fracture features. F o r materials w i t h 0.3 w t . % A l a p r e d o m i n a n t diffusional m e c h a n i s m was p r o p o s e d , while for materials w i t h 1.5 wt. % A l it acts simult a n e o u s l y w i t h a cracking p r o c e s s at low temperatures and w i t h a cavitational m e c h a n i s m at higher temperatures.

1. INTRODUCTION Creep resistance o f t h e r m o m e c h a n i c a l ceramics is o n e o f t h e main p r o p e r t i e s n e e d e d b y designers. F o r b o t h SiC and Si3N4 materials a densification aid is necessary to o b t a i n a high density. T h e creep resistance is strongly a f f e c t e d b y the c h o i c e o f this densification aid material. It has also b e e n s h o w n b y m a n y researchers that, for a given stress and test t e m p e r a t u r e , the creep rate can vary by over several orders o f m a g n i t u d e (see for e x a m p l e refs. 1-3}. With b o r o n or BaC aids, g o o d creep-resistant materials are o b t a i n e d , b u t t h e y still c o n t a i n some p o r o s i t y which lowers their f r a c t u r e resistance and toughness [4, 5]. With o x i d e aids or by sintering u n d e r oxidizing conditions, fully dense materials are o b t a i n e d . A viscous phase is generally p r e s e n t at the grain b o u n d a r i e s and at the triple points. It is thus *Paper presented at the International Symposium on Engineering Ceramics, Jerusalem, Israel, December 16-20, 1984. 0025-5416/85/$3.30

observed t h a t sintering aids play a m o s t imp o r t a n t role in the high t e m p e r a t u r e deform a t i o n and in the creep behaviour, as the carbide (or nitride) crystals themselves do n o t d e f o r m . D e f o r m a t i o n induces cavitation at the interface and triple-point boundaries. T h e effects o f these cavitation processes o n the creep p a r a m e t e r s have been c o n s i d e r e d t h e o r e t i c a l l y by several researchers [ 6 - 1 3 ] b u t t h e r e is a d e a r t h o f e x p e r i m e n t a l data. The aim of the w o r k r e p o r t e d in this p a p e r is to measure creep rates on d i f f e r e n t batches o f hot-pressed SiC materials c o n t a i n i n g different a m o u n t s of a l u m i n i u m and to investigate the influence of m i c r o s t r u c t u r a l p a r a m e t e r s (the a m o u n t and l o c a t i o n o f the s e c o n d a r y phase, and the m e a n grain size o f SiC crystals) on the creep behaviour.

2. EXPERIMENTAL DETAILS 2.1. Materials T h e materials investigated in this p a p e r were hot-pressed SiC o b t a i n e d f r o m the E l e k t r o s c h m e l z w e r k C o m p a n y at K e m p t e n , F.R.G. T h e materials were p r e p a r e d f r o m a-SiC p o w d e r s (with some a l u m i n i u m o r A1203 p o w d e r s as the densification aid) at 2200 °C and at a pressure o f a b o u t 400 MPa. F o u r batches were p r e p a r e d : SiC-0.3wt.%A1, SiC-1.5wt.%A1 and t w o m o r e SiC-1.5wt.%A1 batches which had been subjected to exaggerated g r o w t h ; these batches were designated A03, A15, A15-E1 and A 1 5 - E 2 respectively. T h e materials were received in the f o r m o f discs 75 m m in d i a m e t e r and 3 - 5 m m in thickness. T h e y were g r o u n d to a thickness o f 2 m m and cut into specimens o f size 2 m m × 4 m m × 20 mm. Each specimen was carefully polished with d i a m o n d paste d o w n to 1 p m

© Elsevier Sequoia/Printed in The Netherlands

148

grade and the edges were slightly chamfered. Metallographic investigation of specimens etched by boiling in Murakami's solution was carried out.

gerated growth. The results are presented in Table 1, in which the mean grain size/)e estimated on the assumption of spherical SiC crystals is also given.

2.2. Morphology

2.3. Creep experiments

Transmission electron microscopy showed that all the materials contain a secondary amorphous phase either at some triple points or at triple points and between SiC grains, depending on the a m o u n t of the densification aid (Fig. 1). These regions become increasingly thinner as the distance from the triple points and four-point channels increases. It was not possible, within the resolution of the electron microscope, to show that the glassy phase covers all the grain boundaries. We thus assume that it is essentially confined to triplepoint boundaries and four-point channels [14]. Optical micrographs of A15 and A15-E2 are shown in Fig. 2. The grain size was measured by automatic image analysis using a texture analyser system from Leitz. Between 1200 and 1500 crystals were analysed [15]. Because of the shape of the SiC crystals, we preferred to use the mean surface area parameter -~A rather than the mean grain size of the SiC crystals. The surface area distribution curves for SiC crystals in the different batches are presented in Fig. 3. The results show that the grain size distribution is slightly displaced towards higher values for A15 compared with that for A03 materials. The largest mean grain sizes are observed for SiC materials with exag-

Creep experiments were performed in three-point bending under vacuum (about 10 -6 Torr) in a Vide Moleculaire de l'Industrie Sesame type of opening furnace, with an

Fig. 1. Transmission electron micrograph of SiC0.75wt.%Al, showing the amorphous phase located at triple-point boundaries.

Fig. 2. Optical micrographs of (a) an A15 specimen and (b) an A15-E2 specimen.

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TABLE 1 Morphological characteristics of the SiC-A1 materials Material

-4 A (~ m2 )

De (pm)

A03 A15 A15-E1 A15-E2

3.0 3.2 5.5 9.3

1.0 1.0 1.3 1.7

u p p e r t e m p e r a t u r e limit o f 2 2 7 3 K. The pushrods are m a d e o f t u n g s t e n and were 20 m m in diameter, and the knives are TiC cylinders 5 m m in diameter. The b e n d i n g span is 14 m m . C o n s t a n t loading was achieved b y using a lever s y s t e m and a dead weight acting on the u p p e r push-rod. The l o w e r p u s h - r o d is fixed to the b o t t o m o f the furnace. The u p p e r push-rod is fixed on a cross-head which moves along t w o s u p p o r t columns. Transmission of the load f r o m the lever s y s t e m to the crosshead is via a carbide ball, thus avoiding eccentricity effects. The frictional resistance of this s y s t e m is less t h a n 50 gf [16] (Fig. 4). The d i s p l a c e m e n t is m e a s u r e d on the external part of the push-rod. The signals of t w o s y m m e t r i c a l l y l o c a t e d inductive transducers are a d d e d and p l o t t e d as a f u n c t i o n o f time on a flat-bed recorder. Also t w o 1 / 1 0 0 0 m m dial strain gauges, diametrically o p p o s e d , are used in o r d e r to calibrate the electrical signal

Fig. 4. Creep test equipment.

o f the inductive transducers. A d i s p l a c e m e n t of I p m can easily be measured. Creep e x p e r i m e n t s were u n d e r t a k e n bet w e e n 1673 and 1973 K. The creep temperature was reached within 3 - 4 h using a BRM p r o g r a m m e d electrical p o w e r s u p p l y of 25 k W h . Creep tests were always c o m m e n c e d after a m i n i m u m of I h at the creep-testing t e m p e r a t u r e , to allow for g o o d t e m p e r a t u r e stabilization. The loading s e q u e n c e was always increasing at one t e m p e r a t u r e . S o m e e x p e r i m e n t s were run using o n l y one or t w o loads to c h e c k t h a t in the loading sequence the previous d e f o r m a t i o n has no or o n l y a little influence on the m e a s u r e d strain rates. The creep strain was calculated f r o m the deviation f r o m linearity since all the deform a t i o n s involved are small (less t h a n 1%) [17,181. Scanning e l e c t r o n m i c r o s c o p y using a J E O L T20S i n s t r u m e n t was carried o u t on the specimens in the a s - d e f o r m e d state after slight polishing or polishing and etching. F r a c t o g r a p h y investigations were also m a d e on specimens which b r o k e during creep tests or which were b r o k e n at r o o m t e m p e r a t u r e after creep d e f o r m a t i o n .

150 3. R E S U L T S

All the experimental creep rates were measured in steady state conditions. Such a behaviour appears to exist in these materials as can be seen in Figs. 5 and 6 for A 1 5 - E l materials. The e x t e n t of these linear strain-time (e-t) domains depends on the testing temperature and is probably linked to the ratio o f creep due to grain b o u n d a r y sliding to creep due to cavitation [19]. For Fig. 5 it is difficult to ensure that steady state is n o t a mixture o f primary and tertiary states. Nevertheless, as this domain increases with the increase of t e m p er atu r e as assumed, we considered

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60

80

t (h)

Fig. 5. S t r a i n - t i m e curve for an A15-E1 s p e c i m e n c r e p t at 1673 K and 285 MPa.

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83

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, 10

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, 30

a t (h)

Fig. 6. E x p e r i m e n t a l s t r a i n - t i m e curve for an A15-E1 s p e c i m e n u n d e r successive loadings at 1973 K.

t hat steady state conditions are reached. This choice is confirmed by the fact that the measurement from only one e - t curve under only one load and the use of an incremental m e t h o d as in Fig. 6 give the same value of ~ for this "st eady st at e" stage. To describe the t em perat ure and stress dependence of the creep rate e, we used the following relationship:

where A is the pre-exponential parameter which takes into account the microstructure and the t em perat ure effects, o is the applied stress, Qa is the apparent activation energy, R is the gas constant and T (K) is the temperature. In Fig. 7(a) the stress-strain curves are plotted for A03 and A15 and in Fig. 7(b) for A15, A15-E1 and A15-E2. At 1573 K it was generally not possible to measure any strain rate as the specimens broke either during or shortly after loading or after a little deformation. In all cases the scatter decreases as the t e m p e r a t u r e increases. This is due to the increasing a m o u n t of d e f o r m a t i o n as temperature increases (see Figs. 5 and 6). A general observation is that the A03 specimens exhibit m uch lower strain rates than the A15 specimens. The materials with large grain sizes (A15-E1 and A15-E2) also exhibit lower strain rates than the A15 specimens. If we assume a diffusional mechanism at the grain boundaries, the lower strain rates arise firstly because the a m o u n t of viscous phase is too small to allow free grain b o u n d a r y sliding along the facets and secondly because of the large crystals which impede grain b o u n d a r y sliding as a result of a steric effect. It should also be not ed that two mechanisms are probably acting at 1873 and 1973 K. Table 2 gives the values of the exponential coefficient n at various temperatures. For A03 materials, n (= 1.4) is constant over the whole t em perat ure range investigated while, for all the A15 and A15-E1 materials, n decreases continuously from 2 to 1 when the t e m p e r a t u r e is increased from 1673 to 1973 K. The smaller a m o u n t of experimental data for material A15-E2 prevents accurate measurement of n. Either grain b o u n d a r y embrittlement or very exaggerated growth of some large crystals limits the stress range that

151

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can be investigated at higher values. Figure 7(b) shows that there is only a small difference in the creep rates of the A15-E1 and A15-E2 materials• This is because, although A15-E2 has a larger grain size, the acicular nature of the grain impedes grain boundary sliding• A variation in n with temperature should induce a change in the apparent activation energy QA. This is shown in Fig. 8 for various stress levels. With the A03 and A15-E2 alloys a constant value of 105 kcal mo1-1 is measured• For the A15 and A15-E1 alloys there is a continuous decrease in QA as the temperature decreases; QA depends on both stress and temperature and varies approximately from 30 kcal mo1-1 at 1673 K to 110 kcal mo1-1 at 1973 K. It should be noted that the same value of apparent activation energy is obtained at high temperatures, irrespective of which alloy is considered.

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C ~-- Eve ~- Cdi f -J- Cca v

V a l u e s o f n f o r t h e v a r i o u s a l l o y s in t h e 1 6 7 3 K temperature range

Material

A03 A15 A15-E1

The creep strength of thermomechanical ceramics is strongly dependent on the chemical composition. For SiC-A1 materials the degradation is due to the presence of a secondary viscous phase at the grain boundaries. Lange et al. [20] have suggested that for Si3N 4 materials, several competitive mechanisms occur, often simultaneously, during the creep of thermomechanical ceramics. They are essentially a viscoelastic creep mechanism, a diffusional creep mechanism and a cavitational creep mechanism. The last two are the persistent modes which account for the unrecoverable creep strain. The total creep strain is therefore due to the superimposition of the components of these different types of mechanism [20]:

1973

n f o r the f o l l o w i n g t e m p e r a t u r e s 1673K

1773K

1873K

1973K

1.4 1.6 2.0

1.4 1.5 1.6

1,4 1,2 1,2

1.4 1.0 1.1

The viscoelastic strain eve is predominant in primary creep• It arises as a consequence of the viscoelastic deformation due to grain boundary sliding accommodated by elastic deformation at grain boundary asperities and/ or adjacent grains• In fact, this sliding is hindered by asperities which exist at grain boundaries• The source of these asperities is not obvious and Lange et al. [20] have as-

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Fig. 8. Arrhenius plots for the four SiC-A1 materials at various stresses (as indicated by the numerals on the lines): (a) A03; (b) A15; (c) A15-E1; (d) A15-E2.

sumed that the size of the asperities is approximately equal to the thickness of the glassy phase between the grains (i.e. less than 50 A). The strain •dif is due to the diffusion of one or several constituents of the material. Creep is therefore controlled by the atoms which diffuse most slowly. This deformation is predominant in the steady state. In this case the stress exponent n is equal to unity in the low stress range. The cavitation strain ecav is due to grain boundary sliding. This deformation can be accentuated by subcritical growth of the pre-existing cracks. It predominates in tertiary creep. Hasselman and Venkateswaran [ 10, 11 ] and Venkateswaran and Hasselman [ 13] have proposed that the measured creep strain of

ceramics which show cavitation is due to a superimposition of two terms: the usual viscous creep strain due to grain boundary sliding; the elastic strain which arises from the increase in compliance of the specimen as the material cavitates, i.e. as the number of boundaries between grains decreases. From the relationships proposed for these different mechanisms, it would be possible to calculate the total creep strain from the different creep strain components, supported on the one hand by the experimental creep test re. sults and on the other hand by the fractographic observations on the crept specimens. However, the difficulty is that the stress gradient in three-point bending leads to inhomogeneous cavitation and the relationship,

153 therefore, is very difficult to verify. It should be r e m e m b e r e d that in any case the strain of SiC materials cannot be controlled by any dislocation m o t i o n at temperatures below 1973 K [21, 22]. More recently, Davis e t al. [23] have a t t e m p t e d to show that in the 1 5 7 3 2073 K t e m p e r a t u r e range the creep of SiC materials is controlled by dislocation climb. Their measurements were made on materials which had a high dislocation density in the as-received state. Our materials were almost free of dislocations and no dislocation structure appears after deformation. We observed some dislocation in the crystals but no dislocation structure nor slip bands (Fig. 9). If we analyse the values of the apparent activation energy, our values compare favourably with the values of 110 kcal mo1-1 and 3 0 - 3 5 kcal mo1-1 r epor t e d by Tajima and Kingery [24, 25] in the 2 0 7 3 - 2 4 7 3 K temperature range for the diffusion of aluminium

Fig. 9. Transmission electron micrographs of an A15 specimen after creep deformation at 1733 K and 110 MPa, showing some dislocations.

in SiC and the segregation of aluminium at the grain boundaries respectively. Our values are slightly higher than the value (90 kcal mo1-1) measured by Schnurer e t al. [26] and Grathwohl et al. [27] on similar hot-pressed materials (with 0.25 wt.% A1 and 0.75 wt.% A1) in the same temperature range. T h e y measured a stress e x p o n e n t n of unity, which is lower than the n value (about 1.4) that we measured for A03 alloys. Their experiments were carried out using four-point bending (which may be more accurate than ours in which three-point bending was employed) and at a lower t em perat ure and for smaller stresses than ours. These workers concluded that the creep mechanism is controlled by the diffusion of carbon, on the basis of the diffusion measurements of Hon and coworkers [28-30]. We can conclude that for A03 materials the creep is controlled mainly by a diffusional mechanism with a small influence exerted by a cavitation process. For A15 materials a high degree of grain separation and of cavitation occurs at temperatures higher than 1673 K. The total strain is therefore due, more specifically, to the diffusional and cavitational components. Different mechanisms depending on the test temperature act on ecav. (1) At low temperatures a crack extension c o m p o n e n t occurs which leads to mechanical separation of the crystal facets. In this case there is degradation of the elastic properties of the materials (e.g. the elastic modulus E). (2) At high temperatures a cavitation mechanism takes place by coalescence and growth of the cracks. This mechanism is thermally activated and also leads to a decrease in E. It is interesting to analyse the change in the stress e x p o n e n t n as a funct i on of the temperature (Fig. 10). At a t e m p e r a t u r e below approximately 1823 K, the n values increase as the mean grain size increases. At high temperatures, for the A15 and A15-E1 materials the values are of the same order of magnitude. This probably indicates a morphological change in the d e f o r m a t i o n pattern of these materials. This is confirmed by observation of the deformed and/or fractured materials; for A03 alloys, cavitation appears only at temperatures higher than 1723 K; below this tem-

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1773

1873

1973 T

(K}

Fig. 10. Change in the variation in the stress exponent n with the test temperature for various SiC-AI materials: A, A03; e, A15; ©, A15-E1.

Fig. 11. Micrograph of an A03 spec}men after slight polishing and etching after deformation at 1773 K and which had been strained at 150 MPa (strain rate, about 2 × 10-8 s-l).

the s e c o n d phase is thin e n o u g h to avoid n u c l e a t i o n of spherical cavities at the triple point. In o u r case the density o f cavities is m u c h higher f o r A15 materials t h a n for A03 materials. A t t e m p e r a t u r e s higher t h a n 1723 K the a m o u n t o f facet separation decreases and cavitation is m o r e generalized until 1 9 2 3 - 2 0 7 3 K. F o r A15-E1 and A15-E2 materials, cavitation is also observed and, at lower t e m p e r a t u r e s , intergranular and transgranular r u p t u r e (Fig. 13). It should be n o t e d t h a t the surface o f some grains is highly " g r a n u l a r " . This p o i n t has n o t been clarified; perhaps it is due to a m o r p h o l o g i c a l modification o f the glassy phase a n d / o r an o x i d a t i o n p h e n o m e n o n on the f r a c t u r e surface. T h e observations at l o w e r t e m p e r a t u r e s s h o w for large-crystal materials a very i m p o r t a n t facet separation, which is also p r o b a b l y a c c e n t u a t e d b y the soft etching o f the surface (Fig. 14). F r o m these observations and o u r experim e n t a l results it can be c o n c l u d e d t h a t the m e a n grain size o f the SiC crystals and the a m o u n t o f glassy phase are the p r e d o m i n a n t m o r p h o l o g i c a l p a r a m e t e r s with respect to the creep behaviour. It n o w seems necessary to measure a c c u r a t e l y the change in the p o r o s i t y o f these cavitated materials and p r o b a b l y to use q u a n t i t a t i v e f r a c t o g r a p h y [32] to o b t a i n m o r e i n f o r m a t i o n on the facet separation. F o r example, we can evaluate the a m o u n t o f cavitation, i.e. the loss o f stiffness, which is necessary to explain the curved f e a t u r e o f the A r r h e n i u s plots for the A 1 5 and A15-E1 materials. A c c o r d i n g to the w o r k o f W e e r t m a n [33] and o f Hasselman and V e n k a t e s w a r a n [ 10] the loss o f stiffness d u e to the presence o f cavities and o f m i c r o c r a c k s is o f the f o r m 16

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E = E o ,\ 1 + perature, m i c r o c r a c k s and regions o f glassy phase are the main features observed (Fig. 11). F o r A 1 5 alloys, in the 1 7 2 3 - 1 9 2 3 K t e m p e r a t u r e range, cavitation is very imp o r t a n t ; at l o w e r t e m p e r a t u r e s , grain separation and cavitation can be observed; cavitation ( s o m e t i m e s oblate) e i t h e r is observed at triple-point b o u n d a r i e s or arises t h r o u g h separation o f grain facets (Fig. 12). Marion e t al. [31] c o n s i d e r e d t h a t the n u c l e a t i o n and coalescence o f these oblate cavities arise w h e n

where E 0 is the elastic m o d u l u s in the unc r a c k e d specimen, E the elastic m o d u l u s in the d a m a g e d specimen, N the n u m b e r o f cracks per u n i t v o l u m e and a the size o f the m i c r o c r a c k s or o f the cavities. At a given creep stress the s u p p l e m e n t a r y elastic d i s p l a c e m e n t is thus AG ---- c --Co o 16 Eo 9

Na 3

155

Fig. 12. Micrographs of A15 specimens after slight polishing and etching of the tensile face after deformation at various stresses and temperatures: (a) 1573 K, 416 MPa; (b) 1773 K, 150 MPa (strain rate, about 2 × 10.8 s-l); (c) 1873 K, 260 MPa; (d) 1973 K, deformed up to 20%.

Deriving with respect to time gives the s u p p l e m e n t a r y creep rate as Ae -

48 o N a 2 da 16 o a 3 d_ _N -- + --_ 9 Eo

Fig. 13. Fractography of an A15-E1 specimen deformed at 1873 K under a strain of 160 MPa.

dt

9 Eo

dt

The first t e r m o n the right-hand side describes the influence o f cavity g r o w t h . Irrespective of the m e c h a n i s m , it can be assumed t h a t cavitation is t h e r m a l l y activated. The s e c o n d t e r m describes the increase in microcracks o f size a w h i c h a p p e a r suddenly. We shall call this the m e c h a n i c a l c o n t r i b u t i o n and assume t h a t it is athermal. This m a y be justified for several reasons: the r u p t u r e up to 1 6 7 3 - 1 7 7 3 K is still brittle; the d i s t r i b u t i o n o f the glassy phase b e t w e e n the grains is inh o m o g e n e o u s (there is o f t e n very little glassy

156 For a material with a mean grain size of 1 #m, this represents a microcrack rate of 10 .5 s-1. If we look at the difference between the measured creep rate at 1673 K and the extrapolation on the Arrhenius plots from high temperatures, it can be seen that the value of 10 -s s-1 is of that order of magnitude. So, at lower temperatures, microcracking can explain the low values of the apparent activation energy. As the t e m p e r a t u r e increases, the a m o u n t of microcracking will decrease, whereas thermally activated cavitation will increase. F u r t h e r m o r e the relative a m o u n t of cavitation creep will also decrease. Then the Arrhenius plots will not be so curved irrespective of whether the activation energy represents one or more processes.

5. CONCLUSION

Fig. 14. Micrographs of an A15-E2 specimen (a) before and (b) after deformation at a strain of 270 MPa at 1673 K.

Creep experiments u n d e r t a k e n in the 1 6 7 3 - 1 9 7 3 K t em perat ure range for SiC alloys with different amounts of glassy phase and different crystal sizes have shown that these morphological parameters are predominant. For the A03 materials, the main mechanism is diffusional while, for the A15 materials, it acts simultaneously with a cracking process a~ low temperatures and with a cavitational mechanism at high temperatures. Using the hypotheses of Weertman [33] and Hasselman and Venkateswaran [10] we have shown that it was possible to evaluate the a m o u n t of cavitation and t hat the value obtained is of the same order of magnitude as our experimental results.

ACKNOWLEDGMENTS phase between grain facets and thus a classical cavitation process is unlikely); such facets may thus break in a brittle way because of the local stress concentration, the grain rearrangem e n t and the incapacity of such boundaries to a c c o m m o d a t e some imposed deformation. Thus for a o/Eo of 10 -3, one microcrack per crystal, a microcrack size of 1 pm and a supp l e m en tar y creep rate of 10 -s s-1, dN

3 -

dt

5

×

1013s-1

The authors would like to thank Redouane Moussa (Morocco) who has u n d e r t a k e n all the creep experiments and scanning electron microscope observations for his Th~se de D o c t o r a t ~s Sciences, and Doctors G~rard N o u e t and Jean Vicens for the transmission electron observations. We would also like to thank the Elektroschmelzwerk Company, F.R.G., who kindly provided the SiC materials. This research was supported by the Centre National de la Recherche Scientifique (ATP 9-83-87).

157 REFERENCES

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