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Superplastic behavior of low-doped silicon nitride
MATERIAIS SCIENCE & ‘EWGINEERIHG
ELSEVIER
Materials Science and Engineering A222 (1997) 175-151
Superplastic behavior of low-doped Philippe Laboratoire
de Structure
et PropriJtCs
Burger l, Richard de I’Etat
Solide, URA-CNRS 59655 Wleneuce-d’hcq
silicon nitride
Duclos, Jacques Crampon* 233, BBt. C6, Ukversit& Cede.u, Frame
des Sciences
et Techzologies
de Lille,
Received 29 May 1996; revised 16 August 1996
Abstract Fully dense, fine-grained (0.55 urn) and equiaxed silicon nitride has been produced by hot isostatic pressing performed at 1740°C under 160 MPa. The starting cc-Si,N, powder was mixed with 0.5 wt.% Y,O, and 0.5 wt.% Al,O, used as sintering additives. Prior to deformation a very small amount of glassy phase was detected by transmission electron microscopy at the two-grain and three-grain junctions. In order to study the superplastic deformation of this material, compression tests at constant crosshead rates were conducted up to about 0.5 strain, in nitrogen, from 1600 to 1700°C. The mean stress exponent of the power law was determined to be n = 1 below a transition stress G.* = 20 MPa and II = 0.5 above this stress. Deformation proceeded via diffusion controlled grain boundary sliding with an apparent activation ener,y Q = 614 kJ mol- ‘. The observation of strain whorls, the shear-thickening phenomenon and the transition from a mild to a strong strain hardening, at high stresses, were attributed to the occurrence of rigid and dry grain-to-grain contacts under the dewetting of the intergranular liquid film above the previous transition stress. Keywords:
Ceramics; Silicon nitride; Superplastic deformation
1. Introduction
Many covalent ceramics based on f?ne-grained silicon nitride alloys could be superplastically deformed in the range 1500-1600°C [l-9]. Whatever was the approach used to achieve large strains, all the materials were polyphase ceramics containing a rather large amount of amorphous phase, resulting from the liquid phase sintering and being liquidlike at the elevated test temperatures. This amorphous phase, present either as a nanometer continuous film at the two-grain junctions or as more or less large pockets at multiple-grain junctions [lo-171, has a strong influence on the superplastic flow behavior of the silicon nitride [1,3-81. Thus, it is very important for the development of superplastic S&N, ceramics to reduce the amount of the glass phase up to a lower limit that still allows superplasticity [2]. Hot isostatic pressing (HIPing) is very effective to sinter S&N, powder without or with a very small amount of additives. For example, Homma et al. [18] * Corresponding author. ’ This work is partly based on the Ph.D. Thesis of P. Burger. 0921-5093/97/$17.00 0 1997 Elsevier Science S.A. All rights reserved P1rs0921-5093(96)10519-0
obtained fully dense Si,N, by HIPing pure Si,N, powder without additives at lS50-1950°C under a pressure of 150 MPa. Themelin [19], has HIPed Si,N, powder with the aid of 0.5 wt.% Y,O,. Transmission electron microscopy (TEM) observations have shown a regular distribution of the amorphous phase at the two-grain junctions, the small excess of glassy phase being accumulated at triple-point junctions. Silicon nitride obtained by HIPing S&N, powder with a very small amount of sintering aids as Y203 and Al,O, is concerned by the present work. Intensive grain boundary characterization of such superplastically deformed Si,N, has been previously reported [20]. Deformation is influenced by the interaction between the initially wetting intergranular liquid phase and the S&N, grains, and proceeds via grain boundary sliding accommodated by solution-diffusion-precipitation. A shear-thickening behavior has been observed during compression tests with a stress esponent decreasing with increasing stress from unity to about 0.5 at a transition stress G.* = 20 MPa. This transition is correlated with a mild strain hardening at stresses below G*, which increases strongly above g*. This could not be related to microstructural coarsening alone and arises
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wetting
2. Experimental procedure 2.1. Material
The investigated material was a developmental HIPed silicon nitride, supplied by the Ceramique Technique Desmarquest Lab. (Evreux, France). The starting fine a-Si,N, powder (HC Stark LC 12 S) was mixed with 0.5 wt.% Y,O, and 0.5 wt.% A&O, as sintering aids, and then attrition ball milled in alcohol for 12 h, using silicon nitride milling media. The slurry was dried by heat centrifugation to obtain a deagglomerated powder mixture. Green compacts, 20 mm in diameter and 45 mm height, were prepared by cold isostatic pressing (CIP) at 150 MPa. Rods blanks reaching a density of about 65% of the theoretical density were obtained after an annealing treatment at 1200°C in a nitrogen atmosphere. Finally, rod blanks were fully densified at the ONERA lab (Chatillon, France), by tantalum encapsulated HIPing under an average argon pressure of 160 MPa at 1740°C during 90 min. The samples were then quickly cooled (15’C min-‘) to limit the grain growth. After HIPing, the specimens were ground and polished to 1 urn finish and their density was measured using Archimedes’ principle with methanol as the immersion medium and a sapphire crystal as reference sample. The density of all specimens used for compression tests was the theoretical (3.15 g cmB3). After HIPing the total oxygen content of the specimens was about 2 wt.%. By X-ray diffraction (XRD) analysis the E-P phase transformation was revealed nearly complete, the a phase being less than 10% of the total Si,N, and the microstructure was observed to contain traces of Si,N,O. All HIPed Si,N, materials examined by TEM throughout this study revealed a glassy phase. The glassy phase was rich in cations such as Si, Y, and Al. Based on these evidences, the composition of the glassy phase was assumed to be the same than the eutectic composition in the SiO*-Y,03-Al,O, system. From the above hypothesis and assuming that all the yttria was consumed in the glassy phase, the volume fraction V of the liquidlike phase was calculated to be 1.1% at the eutectic temperature [21]. The excess of alumina and silica was assumed to lead respectively to ,0’-S&N4 and Si,N20 grains as observed. The geometrically determined volume fraction of the liquid phase wetting the grains with a film thickness h, is V= 3/2,/d. For a mean grain size d = 0.55 urn, our calculated volume fraction corresponds to a mean film thickness h, of about 2 nm, which is very closed to the grain boundary film thickness in this class of silicon nitride.
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2.2. Mechanical testing
Uniaxial compression was conducted between 1600 and 1700°C in the stress range 8-150 MPa, under a small overpressure (30 hPa) of flowing high purity nitrogen. Deformation tests were performed using a testing machine (Instron) monitored at constant displacement rates of the mobile crosshead corresponding to initial strain rates in the range 2 x 10e6 to 2 x 10m5 s -1 . Specimens for mechanical tests were cut from densified blanks and ground in 3 x 3 x 7 mm3 bars. The specimens were placed between two Sic disks. A spray of BN was deposited onto the buttons and generally samples were compressed to a maximum true strain equal or lower than 0.5 to avoid friction conditions [3]. To further limit the weight loss generally observed [I] owing to the decomposition of Si3N4 and vaporization of SiO [22] from the sample, a porous silicon nitride ring, fabricated from the green rods, was put around the sample during the tests to serve as a buffer. With this precaution, the weight loss in optimum conditions was reduced to about 0.5% for the sample, while it was 20% for the ring. At the end of certain compression tests, samples were cooled down under load to freeze the pertinent high temperature microstructural features. 2.3. Thermomechanical analysis
The thermomechanical analysis has been made according to the phenomenological high temperature deformation power law: 6 = A# with A = A, exp - QJRT
(2)
where 6 is the strain rate, A, a preexponential structural factor, 0 the applied stress, n the stress exponent, Q the deformation activation energy, R and T having their usual meaning. In a rheological approach a stress exponent n = 1 of the power law corresponds to flow behavior which is formally analogous to Newtonian flow (B= G/V) with a constant apparent viscosity q = (de/ di), while ti > 1 or n < 1 correspond respectively to shear-thinning or to shear-thickening flow. Although deformation tests were performed at constant displacement rates of the crosshead and not at constant true strain rates, the analysis of deformation curves has been made as in the case of constant true strain rate tests. The load and displacement measures were converted into true stress, true strain and true strain rate data. The true strain E was defined as ln(l/l,) where 1 and 1, are respectively the instantaneous and the initial lengths of the specimens, the true stress g and
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o,(l/Z,,) with go the nominal stress and B= $(1,/1) with & the initial strain rate. These three variables can be de&ed at a given moment as functions of the strain and of the applied load only; this justifies the analysis method. The precaution that must be taken in the thermomechanical analysis consists in comparing nearly identical microstructures. Isostructural stress exponents n were deduced by performing strain rate changes during tests at constant temperature. Additional constant stress experiments were performed at 20 MPa in order to determine the activation energy Q, using temperature jumps.
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A
1.910-5 1.210-5
•I 0
7.210-6 2.410-6
q
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2.4. Transmission electron microscopy
Thin foils used for TEM observations were prepared from longitudinal slices, parallel to the compression axis in the case of deformed samples, by mechanically polishing to a thickness of 30 urn. Final thinning was obtained by argon ion milling. Grain sizes (CT) were determined from TEM micrographs using the relation d= 1.56L, where L was the average linear intercept. 3. Results 3.1. Mechanical test
Typical compression tests (G---E) data showing the temperature influence on the shape of the deformation curves obtained at an initial strain rate & of 1.2 x IO- 5 S -l are presented in Fig. 1. Beyond the initial transient 180
,
160
80 60
STRAIN
Fig. 2. Compressive stress-strain curves at 1643°C and at various initial strain rates.
accounting principally for the elastic response of the sample together with the load frame, the stress is always increasing with strain. As the increase in true strain rate as the sample length decreased is the same for all the samples (same initial length I,), the observed variation in the apparent strain hardening is a result of microstructural changes. At 1643°C a set of stress-strain curves at various initial strain rates is shown in Fig. 2. Except at the resulting lower nominal stresses, the flow stress is highly increasing with strain. At intermediate stresses, there is a change from a mild to a more severe observed strain hardening with increasing stress. This situation will be discussed later. The isostructural stress exponent n has been determined by performing strain rate changes at constant temperature 1643°C. The resulting stress jumps, corresponding to stresses in the range 8-100 MPa, yield distinct values for the stress exponent on both sides of a transition stress G* lying near 20 MPa. It has a value of about II = 1 below the transition stress g*, and of about IZ= 0.5 above the transition. In Fig. 3 is a plot of the true strain rate 2 (at a true strain E= 0.07) vs. reciprocal temperature at 20 MPa. The slope of the straight line corresponds to a value of the apparent activation energy equal to 614 kJ mol-‘. 3.2. Mic~ostwctural obsemations
STRAIN
Fig. 1. Compressive stress-strain curves at an initial strain rate 1.2 IO-’ s-’ and at various temperatures.
Typical micrographs of the as-HIPed microstructure are presented in Fig. 4. The average grain size of the as-HIPed material, predominantly ,L?‘-S&N,, was about
17s
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0.55 urn. The fine equiaxed grains were embedded in a secondary glass phase and appeared devoid of any dislocation. The occurrence of a continuous amorphous thin lilm (about 1 nm) along the grain boundaries in the asHIPed samples was confirmed by high resolution TEM
PQ
Examination of the samples by TEM after deformation reveals a coarsening of the microstructure. For instance, as shown in Fig. 5, the mean grain size increased to about 0.87 urn after a strain of 0.4 at 1643°C. However, the grains remain equiaxed and devoid of any dislocation at true strains of about 0.4. After deformation, the micrographs show the absence of cavity nucleation and growth in samples having kept their theoretical density. The TEM in Fig. 5 presents also the most additional features found after deformation in cooled under load samples. They were strain whorls at numerous grain boundaries indicating the presence of highly localized stressed contacts at grain boundaries. The observation of a large number of strain whorls shows that they are found preferentially on those grain boundaries that experience an average compressive stress higher than a critical stress gc of about 30 MPa [20]. 4. Discussion
Due to the low amount of additives all HIPed Si,N, materials examined by TEM throughout this study revealed a material with a quasi-ideal distribution of the glassy phase as a continuous ultrathin tilm at two-grain junctions and very small excess at the triplepoint junctions of fine-grained and equiaxed microstructure. This permits superplastic deformation
1o-sl 10
Fig. 4. TEM micrographs of as-HIPed S&N, showing (scale bar = 0.5 Km): (a) a fine and equiaxed microstructure; (b) grains without dislocations and with small glassy phase pockets (arrows).
-6
5.2
5.3 104/T
5.4 (K -' )
Fig. 3. Temperature dependence of the strain rate.
under stresses below the cavitational stress needed to nucleate triple-point junction cavities [23]. In the more conventional class of hot-pressed or sintered silicon nitride, as a result of larger glassy pockets, larger grain sizes and higher grain aspect ratios, and due to the observed degradation of mechanical properties at elevated temperature, substantial investigations [24-321 have been directed at the analysis of the deformation behavior in terms of cavity nucleation and growth as being an important resulting mechanism of unaccommodated grain boundary sliding.
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In our material, the absence of dislocation inside the grains, the remaining equiaxed grain shape and the absence of cavitation after compressive true strain as large as 0.5, are in agreement with accommodated grain boundary sliding through stochastic neighboring grain intercalation process [33] as the dominant flow mechanism. These observations together with a stress exponent n equal to one at stresses below a transition stress g* = 20 MPa and the grain size dependence of the flow stress, which is partly responsible for the strain hardening seen in Fig. 2 at the lowest stresses, are common among fine-grained materials with an intergranular liquid phase that deform via superplasticity accommodated by solution-diffusion-precipitation [34-361 as believed to be operational here. From previous studies on the plastic deformation of Y,O,/Al,O,-doped silicon nitride it appears that in the domain where accommodated grain boundary sliding exerts the major control of the strain rate, the solutiondiffusion-precipitation mechanism must be controlled by the diffusion of matter through the glass rather than by the transfer of atoms across the glass/crystal interface [25,37]. In that case, the strain rate is related [34] to the viscosity of the liquid phase. The present experimental activation energy value Q = 614 kJ mol - ’ is near to other Q values reported in the deformation of Y,O,/Al,,O,-doped silicon nitride materials [31,32,38,39] for grain boundary sliding process with the viscosity of the glassy phase as the rate-controlling step of the accommodation mechanism.
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The increase in the apparent strain hardening observed in our material at the highest stresses cannot be explained by grain growth alone. It should be correlated to the change in the stress exponent from ?I = 1 at stresses below the transition stress g* = 20 MPa to IZ< 1 at stresses above the transition. The average stress exponent II = 0.5 determined in the stress range 50-100 MPa corresponds in the rheological approach to a shear-thickening flow regime with an increasing apparent viscosity as the strain rate increases. Comparable values of the stress exponent and of the transition stress have been reported for the first time by Chen and Hwang [3] in fine grained silicon nitride tested in compression. A continuum mechanics model, which takes into account the stress variation among differently oriented grain boundaries, has been formulated by these authors to explain the transition. They postulate that, above a transition stress G* (applied compressive stress), a volume fraction of material, increasing with stress and containing ‘rigid and dry’ contacts between grain boundaries, develop in the silicon nitride. The observation of numerous strain whorls in our superplastically strained samples is in favor of the development of such localized rigid grain-to-grain contact points as the stress increases. Above the transition stress the behavior of the material was analyzed by Chen and Hwang [3] using a two‘phases’ model. In this model, the volume fraction of Si,N, material containing ‘rigid and dry’ contacts between grain boundaries was treated as a second ‘phase of hard inclusions’ in a ‘soft matrix’ of the composite. From this model [3] it is possible to determine an apparent stress exponent defined as: TZ= (d In i)/(d In n)
(3) and then a mean stress exponent (II). With the use of an adimensional variable u = G/G.*, we found rz= (u0.75)/(21+ 3) in the shear-thickening domain. In our experimental determination range of the stress exponent (ur = 2.5, u-, = 5) the value of the mean stress exponent for this regime is then:
Qz)= l - 3.75 (Z12 - Zll)ln[(3
Fig. 5. TEM micrograph of HIPed Si,N, showing a coarsening of the microstructure and numerous strain whorls (arrows) after deformation (scale bar = 1 pm).
+ uZ)/(3 + u,)] = 0.44
(4)
which is very close to our measured ~z value of 0.5. The required condition for solution-diffusion-precipitation compressive deformation [34-361 is that grain boundaries wetted by a presumably liquidlike film can support compressive stresses and normal stress gradients. According to the Clarke’s model [40] when a stable thin glassy film is present at a stress-free grain boundary, this means that the equilibrium thickness is determined by the balance between two competing forces, an attractive van der Walls dispersion interaction between the grains on either side of the boundary acting to thin tilm and a repulsivesteric one, due to the
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structure of the intergranular liquid. This model shows that the thin liquid film itself can support compressive stresses up to some extent. By incorporating the concept of Stern layers at the liquid/crystal interfaces Chen and Hwang [3] have attributed the cause for shearthickening to the possible contact of these residual Stern layers when the liquid film present at a stressed grain boundary will become unstable and slightly depleted. Inserting the appropriate numerical values for the different parameter in the Clarke’s model a rough calculation was made by Chen and Hwang [3] to estimate the critical stress G, (local compressive stress) required to entirely squeeze out the grain boundary liquid except for the Stern layers. They found a value of 35 MPa in agreement with their model of shear-thickening flow that yields a critical stress gc = 3a*/2 = 30 MPa [3]. This value is also in agreement with the fact that the observed strain whorls are localized on the grain boundaries that experience a local compressive stress higher than critical stress ~~ = 30 MPa [20]. The model of Clarke provides clearly a support for the partial dewetting of the grain boundaries at strain whorl places which increases with increasing applied stress. In addition to the case of superplastic silicon nitride (here and in the early work of Chen and Hwang [3]), it has been also used to explain an apparent work hardening behavior in superplastic zirconia [41,42], where dewetting was observed to occur with a minimum applied compressive stress of 20 MPa. The transition in the apparent strain hardening observed at the transition stress can most likely be attributed to the progressive dewetting of the grain boundaries. It has been shown that the model of shearthickening flow can be used to predict a strain hardening the more important as the stress exponent decreases in the shear-thickening regime [22,43]. If, at a compressive stress below o*, an increase of the apparent viscosity is caused by microstructural coarsening, then at compressive stress above g*, the in situ ‘rigid inclusion reinforcement’ due to grain contacts had a cumulative effect with that of grain coarsening on the strain hardening. 5. Conclusion
Silicon nitride with a mean grain size of 0.55 pm was obtained by HIPing at 1740°C. Due to the low amount of additives (0.5 wt.% Y20, and 0.5 wt.% Al,O,) the equiaxed grains were embedded in a very small amount of intergranular glassy phase. Superplastic deformation of this material proceeds with grain boundary sliding controlled by matter transport via diffusion through the glass. After compression tests up to about a strain of 0.5 in the temperature range 1600-17OO”C, the numerous strain whorls together with the changes in the stress
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exponent and in the strain hardening occurring above 20 MPa are in good agreement with the shear-thickening model proposed by Chen and Hwang [3] for superplastic ‘SiAlON’.
Acknowledgements
The M.R.T. is acknowledged for financial support of this work. We would like to thank B. Gales from C.T.D. and M. Parlier from ONERA for help with the material elaboration.
References
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(1) (1995) 211. UOI D.R. Clarke and G. Thomas, J. Am. Ceram. Sot., 60 (11/12) (1977) 491. 1111L.K.V. Lou, T.E. Mitchell and A.H. Heuer, J. Am. Ceratn. Sot., 61 (S/10) (1978) 392. WI A. Thorel, Interface Atomic Structure in SiAlOn Type Silicon Nitride, Ph.D. T/zesis, Ecole Nationale Mines Paris, France, 1988. [I31 H.J. Kleebe, M.K. Cinibulk, R.M. Cannon and M. Ruble, J. Am. Ceram. Sot., 76 (9) (1993) 1969. [I41 M.K. Cinibulk, H.J. Kleebe, G.A. Schneider and M. Rhule, J. Am. Ceratn. Sot., 76 (10) (1993) 2801. 1151 I. Tanaka, K. Igashira, H.J. Kleebe and M. Ruhle, J. .4m. Cerattz. Sot., 77 (1) (1994) 275. [I61 I. Tanka, H.J. Kleebe, M.K. Cinibulk, J. Burley, D.R. Clarke and M. Ruhle, J. Am. Ceram. Sot., 77 (3) (1994) 911. R. Duclos, J. Crampton and B. Cales, [I71 G. Bernard-Granger, submitted for publication in J. Am. Ceram. Sot. [181K. Homma, H. Okada, T. Fujikama and T. Tatuno, J. Ceram. Sot. Jpn. Int. Ed., 95 (1987) 195. [191L. Themelin, Mechanical Behavior and Oxidation Resistance of Silicon and Titanium Nitrides Sintered or Post-Sintered by Hot Isostatic Pressure, Doctorat Tllesis, University of Limoges, France, 1989. VOI P. Burger, R. Duclos and J. Crampon, submitted for publication in J. Am. Ceram. Sot. [211P. Burger, High Temperature Plasticity and Deformation Microstructure of Y,O,-Al,O,-Based HIPed Silicon Nitride, Doctorat Thesis, University of Lille 1, France, 1994. F7-1F.F. Lange, J. Am. Ceratn. Sot. (1982) C-120. Am.
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[35] F. Wakai, Acta Metall. Mater., 42 (4) (1994) 1163. [36] F. Wakai, N. Kondo, H. Ogawa.and H.J. Kleebe, Proc. 4th Euro Ceramics, Riccione, Italy, October 1995. Faenza Editrice S.p.A., Faenza, Italy, 1995, p. 501. [37] J. Crampon, R. Duclos and N. Rakotoharisoa, J. Mater. Sci., 25 (1993) 909. [38] L.J. Bowen, T.J. Carruthers and R.J. Brook, J. Am. Ceranz. SOL, 61 (1978) 335. [39] F. Peni, J. Crampon and R. Duclos, Ceram. ht., 18 (1992) 413. [40] D.R. Clarke, J. Am. Ceram. SOL, 70 (1) (1987) 15. [41] M.M.R. Boutz, A.J.A. Winnubst and A.J. Burggraaf, J. Am Ceram. SOL, 77 (10) (1994) 2632. [42] M.M.R. Boutz, A.J.A. Winnubst, A.J. Burggraaf, M. Nauer and C. Carry, J. Eur. Cernm. Sot., (13) (1994) 103. [43] P. Burger, F. Peni, J. Crampon, R. Duclos, B. Cales and M. Parlier, Proc. 3rd Euro Ceramics, Madrid, Spain, Septerrlbet 1993, Faenza Editrice Iberica S.L., Castellon de la Plana, Spain, 1993, p. 809.
ELSEVIER
Materials Science and Engineering A222 (1997) 175-151
Superplastic behavior of low-doped Philippe Laboratoire
de Structure
et PropriJtCs
Burger l, Richard de I’Etat
Solide, URA-CNRS 59655 Wleneuce-d’hcq
silicon nitride
Duclos, Jacques Crampon* 233, BBt. C6, Ukversit& Cede.u, Frame
des Sciences
et Techzologies
de Lille,
Received 29 May 1996; revised 16 August 1996
Abstract Fully dense, fine-grained (0.55 urn) and equiaxed silicon nitride has been produced by hot isostatic pressing performed at 1740°C under 160 MPa. The starting cc-Si,N, powder was mixed with 0.5 wt.% Y,O, and 0.5 wt.% Al,O, used as sintering additives. Prior to deformation a very small amount of glassy phase was detected by transmission electron microscopy at the two-grain and three-grain junctions. In order to study the superplastic deformation of this material, compression tests at constant crosshead rates were conducted up to about 0.5 strain, in nitrogen, from 1600 to 1700°C. The mean stress exponent of the power law was determined to be n = 1 below a transition stress G.* = 20 MPa and II = 0.5 above this stress. Deformation proceeded via diffusion controlled grain boundary sliding with an apparent activation ener,y Q = 614 kJ mol- ‘. The observation of strain whorls, the shear-thickening phenomenon and the transition from a mild to a strong strain hardening, at high stresses, were attributed to the occurrence of rigid and dry grain-to-grain contacts under the dewetting of the intergranular liquid film above the previous transition stress. Keywords:
Ceramics; Silicon nitride; Superplastic deformation
1. Introduction
Many covalent ceramics based on f?ne-grained silicon nitride alloys could be superplastically deformed in the range 1500-1600°C [l-9]. Whatever was the approach used to achieve large strains, all the materials were polyphase ceramics containing a rather large amount of amorphous phase, resulting from the liquid phase sintering and being liquidlike at the elevated test temperatures. This amorphous phase, present either as a nanometer continuous film at the two-grain junctions or as more or less large pockets at multiple-grain junctions [lo-171, has a strong influence on the superplastic flow behavior of the silicon nitride [1,3-81. Thus, it is very important for the development of superplastic S&N, ceramics to reduce the amount of the glass phase up to a lower limit that still allows superplasticity [2]. Hot isostatic pressing (HIPing) is very effective to sinter S&N, powder without or with a very small amount of additives. For example, Homma et al. [18] * Corresponding author. ’ This work is partly based on the Ph.D. Thesis of P. Burger. 0921-5093/97/$17.00 0 1997 Elsevier Science S.A. All rights reserved P1rs0921-5093(96)10519-0
obtained fully dense Si,N, by HIPing pure Si,N, powder without additives at lS50-1950°C under a pressure of 150 MPa. Themelin [19], has HIPed Si,N, powder with the aid of 0.5 wt.% Y,O,. Transmission electron microscopy (TEM) observations have shown a regular distribution of the amorphous phase at the two-grain junctions, the small excess of glassy phase being accumulated at triple-point junctions. Silicon nitride obtained by HIPing S&N, powder with a very small amount of sintering aids as Y203 and Al,O, is concerned by the present work. Intensive grain boundary characterization of such superplastically deformed Si,N, has been previously reported [20]. Deformation is influenced by the interaction between the initially wetting intergranular liquid phase and the S&N, grains, and proceeds via grain boundary sliding accommodated by solution-diffusion-precipitation. A shear-thickening behavior has been observed during compression tests with a stress esponent decreasing with increasing stress from unity to about 0.5 at a transition stress G.* = 20 MPa. This transition is correlated with a mild strain hardening at stresses below G*, which increases strongly above g*. This could not be related to microstructural coarsening alone and arises
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wetting
2. Experimental procedure 2.1. Material
The investigated material was a developmental HIPed silicon nitride, supplied by the Ceramique Technique Desmarquest Lab. (Evreux, France). The starting fine a-Si,N, powder (HC Stark LC 12 S) was mixed with 0.5 wt.% Y,O, and 0.5 wt.% A&O, as sintering aids, and then attrition ball milled in alcohol for 12 h, using silicon nitride milling media. The slurry was dried by heat centrifugation to obtain a deagglomerated powder mixture. Green compacts, 20 mm in diameter and 45 mm height, were prepared by cold isostatic pressing (CIP) at 150 MPa. Rods blanks reaching a density of about 65% of the theoretical density were obtained after an annealing treatment at 1200°C in a nitrogen atmosphere. Finally, rod blanks were fully densified at the ONERA lab (Chatillon, France), by tantalum encapsulated HIPing under an average argon pressure of 160 MPa at 1740°C during 90 min. The samples were then quickly cooled (15’C min-‘) to limit the grain growth. After HIPing, the specimens were ground and polished to 1 urn finish and their density was measured using Archimedes’ principle with methanol as the immersion medium and a sapphire crystal as reference sample. The density of all specimens used for compression tests was the theoretical (3.15 g cmB3). After HIPing the total oxygen content of the specimens was about 2 wt.%. By X-ray diffraction (XRD) analysis the E-P phase transformation was revealed nearly complete, the a phase being less than 10% of the total Si,N, and the microstructure was observed to contain traces of Si,N,O. All HIPed Si,N, materials examined by TEM throughout this study revealed a glassy phase. The glassy phase was rich in cations such as Si, Y, and Al. Based on these evidences, the composition of the glassy phase was assumed to be the same than the eutectic composition in the SiO*-Y,03-Al,O, system. From the above hypothesis and assuming that all the yttria was consumed in the glassy phase, the volume fraction V of the liquidlike phase was calculated to be 1.1% at the eutectic temperature [21]. The excess of alumina and silica was assumed to lead respectively to ,0’-S&N4 and Si,N20 grains as observed. The geometrically determined volume fraction of the liquid phase wetting the grains with a film thickness h, is V= 3/2,/d. For a mean grain size d = 0.55 urn, our calculated volume fraction corresponds to a mean film thickness h, of about 2 nm, which is very closed to the grain boundary film thickness in this class of silicon nitride.
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2.2. Mechanical testing
Uniaxial compression was conducted between 1600 and 1700°C in the stress range 8-150 MPa, under a small overpressure (30 hPa) of flowing high purity nitrogen. Deformation tests were performed using a testing machine (Instron) monitored at constant displacement rates of the mobile crosshead corresponding to initial strain rates in the range 2 x 10e6 to 2 x 10m5 s -1 . Specimens for mechanical tests were cut from densified blanks and ground in 3 x 3 x 7 mm3 bars. The specimens were placed between two Sic disks. A spray of BN was deposited onto the buttons and generally samples were compressed to a maximum true strain equal or lower than 0.5 to avoid friction conditions [3]. To further limit the weight loss generally observed [I] owing to the decomposition of Si3N4 and vaporization of SiO [22] from the sample, a porous silicon nitride ring, fabricated from the green rods, was put around the sample during the tests to serve as a buffer. With this precaution, the weight loss in optimum conditions was reduced to about 0.5% for the sample, while it was 20% for the ring. At the end of certain compression tests, samples were cooled down under load to freeze the pertinent high temperature microstructural features. 2.3. Thermomechanical analysis
The thermomechanical analysis has been made according to the phenomenological high temperature deformation power law: 6 = A# with A = A, exp - QJRT
(2)
where 6 is the strain rate, A, a preexponential structural factor, 0 the applied stress, n the stress exponent, Q the deformation activation energy, R and T having their usual meaning. In a rheological approach a stress exponent n = 1 of the power law corresponds to flow behavior which is formally analogous to Newtonian flow (B= G/V) with a constant apparent viscosity q = (de/ di), while ti > 1 or n < 1 correspond respectively to shear-thinning or to shear-thickening flow. Although deformation tests were performed at constant displacement rates of the crosshead and not at constant true strain rates, the analysis of deformation curves has been made as in the case of constant true strain rate tests. The load and displacement measures were converted into true stress, true strain and true strain rate data. The true strain E was defined as ln(l/l,) where 1 and 1, are respectively the instantaneous and the initial lengths of the specimens, the true stress g and
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o,(l/Z,,) with go the nominal stress and B= $(1,/1) with & the initial strain rate. These three variables can be de&ed at a given moment as functions of the strain and of the applied load only; this justifies the analysis method. The precaution that must be taken in the thermomechanical analysis consists in comparing nearly identical microstructures. Isostructural stress exponents n were deduced by performing strain rate changes during tests at constant temperature. Additional constant stress experiments were performed at 20 MPa in order to determine the activation energy Q, using temperature jumps.
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•I 0
7.210-6 2.410-6
q
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2.4. Transmission electron microscopy
Thin foils used for TEM observations were prepared from longitudinal slices, parallel to the compression axis in the case of deformed samples, by mechanically polishing to a thickness of 30 urn. Final thinning was obtained by argon ion milling. Grain sizes (CT) were determined from TEM micrographs using the relation d= 1.56L, where L was the average linear intercept. 3. Results 3.1. Mechanical test
Typical compression tests (G---E) data showing the temperature influence on the shape of the deformation curves obtained at an initial strain rate & of 1.2 x IO- 5 S -l are presented in Fig. 1. Beyond the initial transient 180
,
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80 60
STRAIN
Fig. 2. Compressive stress-strain curves at 1643°C and at various initial strain rates.
accounting principally for the elastic response of the sample together with the load frame, the stress is always increasing with strain. As the increase in true strain rate as the sample length decreased is the same for all the samples (same initial length I,), the observed variation in the apparent strain hardening is a result of microstructural changes. At 1643°C a set of stress-strain curves at various initial strain rates is shown in Fig. 2. Except at the resulting lower nominal stresses, the flow stress is highly increasing with strain. At intermediate stresses, there is a change from a mild to a more severe observed strain hardening with increasing stress. This situation will be discussed later. The isostructural stress exponent n has been determined by performing strain rate changes at constant temperature 1643°C. The resulting stress jumps, corresponding to stresses in the range 8-100 MPa, yield distinct values for the stress exponent on both sides of a transition stress G* lying near 20 MPa. It has a value of about II = 1 below the transition stress g*, and of about IZ= 0.5 above the transition. In Fig. 3 is a plot of the true strain rate 2 (at a true strain E= 0.07) vs. reciprocal temperature at 20 MPa. The slope of the straight line corresponds to a value of the apparent activation energy equal to 614 kJ mol-‘. 3.2. Mic~ostwctural obsemations
STRAIN
Fig. 1. Compressive stress-strain curves at an initial strain rate 1.2 IO-’ s-’ and at various temperatures.
Typical micrographs of the as-HIPed microstructure are presented in Fig. 4. The average grain size of the as-HIPed material, predominantly ,L?‘-S&N,, was about
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0.55 urn. The fine equiaxed grains were embedded in a secondary glass phase and appeared devoid of any dislocation. The occurrence of a continuous amorphous thin lilm (about 1 nm) along the grain boundaries in the asHIPed samples was confirmed by high resolution TEM
PQ
Examination of the samples by TEM after deformation reveals a coarsening of the microstructure. For instance, as shown in Fig. 5, the mean grain size increased to about 0.87 urn after a strain of 0.4 at 1643°C. However, the grains remain equiaxed and devoid of any dislocation at true strains of about 0.4. After deformation, the micrographs show the absence of cavity nucleation and growth in samples having kept their theoretical density. The TEM in Fig. 5 presents also the most additional features found after deformation in cooled under load samples. They were strain whorls at numerous grain boundaries indicating the presence of highly localized stressed contacts at grain boundaries. The observation of a large number of strain whorls shows that they are found preferentially on those grain boundaries that experience an average compressive stress higher than a critical stress gc of about 30 MPa [20]. 4. Discussion
Due to the low amount of additives all HIPed Si,N, materials examined by TEM throughout this study revealed a material with a quasi-ideal distribution of the glassy phase as a continuous ultrathin tilm at two-grain junctions and very small excess at the triplepoint junctions of fine-grained and equiaxed microstructure. This permits superplastic deformation
1o-sl 10
Fig. 4. TEM micrographs of as-HIPed S&N, showing (scale bar = 0.5 Km): (a) a fine and equiaxed microstructure; (b) grains without dislocations and with small glassy phase pockets (arrows).
-6
5.2
5.3 104/T
5.4 (K -' )
Fig. 3. Temperature dependence of the strain rate.
under stresses below the cavitational stress needed to nucleate triple-point junction cavities [23]. In the more conventional class of hot-pressed or sintered silicon nitride, as a result of larger glassy pockets, larger grain sizes and higher grain aspect ratios, and due to the observed degradation of mechanical properties at elevated temperature, substantial investigations [24-321 have been directed at the analysis of the deformation behavior in terms of cavity nucleation and growth as being an important resulting mechanism of unaccommodated grain boundary sliding.
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In our material, the absence of dislocation inside the grains, the remaining equiaxed grain shape and the absence of cavitation after compressive true strain as large as 0.5, are in agreement with accommodated grain boundary sliding through stochastic neighboring grain intercalation process [33] as the dominant flow mechanism. These observations together with a stress exponent n equal to one at stresses below a transition stress g* = 20 MPa and the grain size dependence of the flow stress, which is partly responsible for the strain hardening seen in Fig. 2 at the lowest stresses, are common among fine-grained materials with an intergranular liquid phase that deform via superplasticity accommodated by solution-diffusion-precipitation [34-361 as believed to be operational here. From previous studies on the plastic deformation of Y,O,/Al,O,-doped silicon nitride it appears that in the domain where accommodated grain boundary sliding exerts the major control of the strain rate, the solutiondiffusion-precipitation mechanism must be controlled by the diffusion of matter through the glass rather than by the transfer of atoms across the glass/crystal interface [25,37]. In that case, the strain rate is related [34] to the viscosity of the liquid phase. The present experimental activation energy value Q = 614 kJ mol - ’ is near to other Q values reported in the deformation of Y,O,/Al,,O,-doped silicon nitride materials [31,32,38,39] for grain boundary sliding process with the viscosity of the glassy phase as the rate-controlling step of the accommodation mechanism.
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The increase in the apparent strain hardening observed in our material at the highest stresses cannot be explained by grain growth alone. It should be correlated to the change in the stress exponent from ?I = 1 at stresses below the transition stress g* = 20 MPa to IZ< 1 at stresses above the transition. The average stress exponent II = 0.5 determined in the stress range 50-100 MPa corresponds in the rheological approach to a shear-thickening flow regime with an increasing apparent viscosity as the strain rate increases. Comparable values of the stress exponent and of the transition stress have been reported for the first time by Chen and Hwang [3] in fine grained silicon nitride tested in compression. A continuum mechanics model, which takes into account the stress variation among differently oriented grain boundaries, has been formulated by these authors to explain the transition. They postulate that, above a transition stress G* (applied compressive stress), a volume fraction of material, increasing with stress and containing ‘rigid and dry’ contacts between grain boundaries, develop in the silicon nitride. The observation of numerous strain whorls in our superplastically strained samples is in favor of the development of such localized rigid grain-to-grain contact points as the stress increases. Above the transition stress the behavior of the material was analyzed by Chen and Hwang [3] using a two‘phases’ model. In this model, the volume fraction of Si,N, material containing ‘rigid and dry’ contacts between grain boundaries was treated as a second ‘phase of hard inclusions’ in a ‘soft matrix’ of the composite. From this model [3] it is possible to determine an apparent stress exponent defined as: TZ= (d In i)/(d In n)
(3) and then a mean stress exponent (II). With the use of an adimensional variable u = G/G.*, we found rz= (u0.75)/(21+ 3) in the shear-thickening domain. In our experimental determination range of the stress exponent (ur = 2.5, u-, = 5) the value of the mean stress exponent for this regime is then:
Qz)= l - 3.75 (Z12 - Zll)ln[(3
Fig. 5. TEM micrograph of HIPed Si,N, showing a coarsening of the microstructure and numerous strain whorls (arrows) after deformation (scale bar = 1 pm).
+ uZ)/(3 + u,)] = 0.44
(4)
which is very close to our measured ~z value of 0.5. The required condition for solution-diffusion-precipitation compressive deformation [34-361 is that grain boundaries wetted by a presumably liquidlike film can support compressive stresses and normal stress gradients. According to the Clarke’s model [40] when a stable thin glassy film is present at a stress-free grain boundary, this means that the equilibrium thickness is determined by the balance between two competing forces, an attractive van der Walls dispersion interaction between the grains on either side of the boundary acting to thin tilm and a repulsivesteric one, due to the
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structure of the intergranular liquid. This model shows that the thin liquid film itself can support compressive stresses up to some extent. By incorporating the concept of Stern layers at the liquid/crystal interfaces Chen and Hwang [3] have attributed the cause for shearthickening to the possible contact of these residual Stern layers when the liquid film present at a stressed grain boundary will become unstable and slightly depleted. Inserting the appropriate numerical values for the different parameter in the Clarke’s model a rough calculation was made by Chen and Hwang [3] to estimate the critical stress G, (local compressive stress) required to entirely squeeze out the grain boundary liquid except for the Stern layers. They found a value of 35 MPa in agreement with their model of shear-thickening flow that yields a critical stress gc = 3a*/2 = 30 MPa [3]. This value is also in agreement with the fact that the observed strain whorls are localized on the grain boundaries that experience a local compressive stress higher than critical stress ~~ = 30 MPa [20]. The model of Clarke provides clearly a support for the partial dewetting of the grain boundaries at strain whorl places which increases with increasing applied stress. In addition to the case of superplastic silicon nitride (here and in the early work of Chen and Hwang [3]), it has been also used to explain an apparent work hardening behavior in superplastic zirconia [41,42], where dewetting was observed to occur with a minimum applied compressive stress of 20 MPa. The transition in the apparent strain hardening observed at the transition stress can most likely be attributed to the progressive dewetting of the grain boundaries. It has been shown that the model of shearthickening flow can be used to predict a strain hardening the more important as the stress exponent decreases in the shear-thickening regime [22,43]. If, at a compressive stress below o*, an increase of the apparent viscosity is caused by microstructural coarsening, then at compressive stress above g*, the in situ ‘rigid inclusion reinforcement’ due to grain contacts had a cumulative effect with that of grain coarsening on the strain hardening. 5. Conclusion
Silicon nitride with a mean grain size of 0.55 pm was obtained by HIPing at 1740°C. Due to the low amount of additives (0.5 wt.% Y20, and 0.5 wt.% Al,O,) the equiaxed grains were embedded in a very small amount of intergranular glassy phase. Superplastic deformation of this material proceeds with grain boundary sliding controlled by matter transport via diffusion through the glass. After compression tests up to about a strain of 0.5 in the temperature range 1600-17OO”C, the numerous strain whorls together with the changes in the stress
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exponent and in the strain hardening occurring above 20 MPa are in good agreement with the shear-thickening model proposed by Chen and Hwang [3] for superplastic ‘SiAlON’.
Acknowledgements
The M.R.T. is acknowledged for financial support of this work. We would like to thank B. Gales from C.T.D. and M. Parlier from ONERA for help with the material elaboration.
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