silicon nitride-silicon carbide composite materials

Refractory Metals & Hard Materials 11 (1992) 213-221

Silicon Nitride-Silicon Carbide Composite Materials

V. Biasini, S. Guicciardi & A. Bellosi IRTEC-CNR, Research Institute for Ceramics Technology, via Granarolo, 64, Faenza, Italy

(Received 21 August 1992; accepted 19 October 1992)

Abstract: Starting from different Si3N 4 matrices, submicrometric 20 vol. % SiC particle-reinforced Si3N 4 composites were obtained by hot-pressing and pressureless sintering. Using the Ray and Bordia model, the sintering behaviour of the Si3N4 matrices during hot-pressing was found to be influenced by the presence of the second-phase particles. Mechanical properties of the matrices and composites were demonstrated to be strongly dependent upon grain size, sintering aid amount, densification route and particle reinforcement.

INTRODUCTION

Silicon nitride and silicon carbide ceramics are leading candidates for high-temperature struct- ural applications, owing to their ability to operate under severe conditions such as high pressure, large temperature fluctuations, acid environ- ments, etc. Unfortunately, their intrinsic brittle- ness is still the main limitation and drawback for such applications.

Previous research on multi-phase ceramics has shown that the mechanical properties of mono- lithic ceramics can be enhanced through the composite approach, including the reinforcement by fibres and whiskers 1-4 and by particle disper- sions. 5-~4 Considering the Si3N4-SiC system, where the reinforcing phase (SIC) is present in the form of fibres or whiskers, a strong improvement in fracture toughness was generally observed, with values up to 12-13 MPa ml/2. 2 In some cases, on the other hand, such second phases were respon- sible for mechanical strength degradation. More- over, whiskers and fibre reinforcement presents technological and reliability problems, besides being an expensive solution. Therefore, increasing attention was devoted to particulate addition and it was reported that the fracture toughness of sift- con nitride could be increased by using SiC particles as dispersoids. ~

Greil et al. 7 found that it was possible to improve the strength in hipped composites even without a parallel toughness enhancement. In the Si3N4-SiC system, the mechanical properties are related to particle size distribution of the SiC powder: Akimune 6 obtained the best results on Si3N4-SiC composites (strength up to 1226 MPa and K~c = 5.4 MPa m L/2) using ultrafine powders and pressureless sintering followed by post hipping.

An important factor in strengthening and toughening mechanism analysis is the relationship between the dimensions of second-phase particles and the matrix grains. 8 If the grain size of SiC is higher than a critical value, both toughness and strength are strongly affected: the former increases and the latter decreases with increase in the reinforcing particle size.

In the present study, SiC-particle/Si3N 4 composites, containing 20 vol. % of SiC (Lonza Carbo- gran UF 25), were produced by hot-pressing and pressureless sintering, starting from two commercial Si3N 4 powders (Starck LC12, UBE SN10).

Reinforcing-phase submicrometer-sized a-SiC particles were chosen because, as reported in literature, 8 in the Si3N4-SiC composites small particles are less deleterious to the bulk strength than large particles.

Refractory Metals &Hard Materials 0263-4368/93 S06.00 © 1993 Elsevier Science Publishers Ltd, England. Printed in Great Britain,

214 V. Biasini, S. Guicciardi, A . Bellosi

Table 1. Characteristics of starting powders

Si~N4 Specific surface Equivalent spherical Aggregate mean a-S(~N 4 area On 2/g) diameter (l~rn) size (i.trn)

Impurities content (%)

0 Ca F AI C C7

StarckLC12 18"8 0-10 0"75 93"0 1.98 0'008 (t'008 0"048 0"16 -- UBESN 10 10'6 0"19 0-70 95.0 1-40 0.005 0"010 0"005 -- 0-01

SiC Specific surface Equivalent spherical Impurities content (%) area (rn :lg) diameter (/~m)

O, Fe204 A LO~ Clree Sire,, ~

Lonza Carbogran UF 25 25-0 0"(/8 1'70 0.05 0'03 0.30 0" 14

Table 2. Composition of the tested mixtures

A SigN 4 Starck LCI2 + 3 wt. % A1203" + 8 wt. % Y203 t~ B SigN4 UBE SN10 + 3 wt. % A I 2 0 / ' + 8 wt. % Y203 h D SigN4 UBE SN10 + 2 wt. % A1203"+ 5 wt. % Y203 b

AC mix A+ 20 vol. % SiC ~ BC mix B + 2 0 v o l . % S i C ' DC mix D + 20 vol. % SiC <

"A120~ Alcoa A16. t'Y~O~ Rhone Poulenc. 'SiC Lonza Carbogram UF25.

EXPERIMENTAL PROCEDURE

Dense specimens were prepared from the starting powders whose characteristics are reported in Table 1. Three silicon nitride matrices, A, B and D, and the corresponding composites AC, BC and DC, were prepared using the compositions shown in Table 2.

Mixes A and B were prepared by powder homogenization in a plastic jar with water and silicon nitride balls as the milling media. They were freeze-dried and then sieved down to 400 ~m. In mix D, the sintering aids were added through a chemical process of co-precipitation of yttria and alumina nitrates on a dispersed Si3N 4 suspension with pH > 10. The slurry was washed and freeze- dried, and the resulting powder was then calcined at 400°C and sieved down to 400/~m.

The SiC powder was added to the previous blends by homogenization in isobuthyl alcohol with silicon nitride balls as milling media. The composite mixtures (AC, BC and DC) were dried in an oven at 150°C and sieved down to 250/.tm.

Densification tests were performed by pressureless sintering in N2 atmosphere at 1800°C for 120 min and by hot-pressing in vacuo in an induc- tion-heated graphite die at temperatures of

1715-1750°C with an applied pressure of 30 MPa.

On the dense samples, microstructural characteristics were investigated by X-ray diffraction, scanning electron microscopy (SEM) and X-ray microanalysis. The Young's modulus (E) was measured on samples with dimensions 0.8 mm x 8"0 mm x 28 mm (height by width by length, respectively) by the frequency resonance method. ~5 Flexural strength, both at room and high temperature, was measured on samples 3 mm x 3 mmx 30 mm in a 4-point bending fixture, with 26 mm outer span, 13 mm inner span, and a crosshead speed of 0.5 mm/min. In the high temperature tests, the specimen was held at the set temperature for 10 rain before loading to ensure thermal equilibrium. The bars had been previously machined lengthwise with a 100-grit resin- bonded diamond wheel. The last 0-1 mm of thickness was removed at 4 ttm/pass. Chamfering of the edges was performed in the same way. The fracture toughness (K~c) was evaluated by the direct crack measurement method with a load of 98 N in a hardness tester Zwick 3112 using the formula proposed by Anstis et al. 16 On the same apparatus, the Vickers microhardness ( H V ) was measured on polished surfaces with a load of 4.91 N.

RESULTS AND DISCUSSION

Densification and microstructure

Figure lia) and (b) shows the fractional density of the samples during the heating up and the subsequent soaking stage under hot-pressing conditions.

Silicon nitride-silicon carbide composite materials 215

Fig. 1.

80

70

60

50

40

A

- - I - - AC

- ~ B ,/

1 t_ L

1000

100

90

80

70

P/

1

Oth(%)

1200 1400 1600

T(°C) ( a :

1800

i I

12 30

t(min) (b)

Relative density of the hot-pressed samples during the heating stage (a) and the subsequent soaking time (b).

With reference to the samples produced by mechanical mixing (samples A, B, AC and BC), the shrinkage started at about 1550°C for the baseline matrices (samples A and B), with an initial green density of 66% for the specimen A and 52% for the specimen B. The related composites, AC and BC, started to densify at 1440°C with green densities of 52% and 60%, respectively.

With reference to the samples produced by chemical co-precipitation of the sintering aids

(sample D and DC), an initial stage, where the shrinkage can be mainly attributed to particle rearrangement, ~7 occurred in the temperature range between 1000 and 1500°C; during this stage the green density rose from 45 to 51%.

For all the composites, no evident densification reduction, imputable to the second-phase presence, was detected during the heating stage up to 1700°C. Slight differences in the sinterability can mainly be ascribed to the baseline matrix characteristics. The maximum densification rates

216 V. Biasini, S. Guicciardi, A. Bellosi

occurred in the range 1600-1700°C. A compari- son among densification behaviors of the samples containing 3 wt % A1203 and 8 wt % Y203 in the stage of constant temperature is only possible between matrixes A and B, which were both hot- pressed at 1710°C, and composites AC and BC, hot-pressed at 1750°C. The B and BC materials produced with UBE silicon nitride showed the best sinterability, reaching full density in a very short time. Samples produced with Stark Si3N4, A and AC, exhibited a slightly lower sinterability, particularly above the relative density of 90%. The samples D and DC, containing the lowest amount of sintering aids (2 wt % AI203 and 5 wt % Y203) , were characterized by slower densification rate and needed longer soaking times to completely densify. In any case, all the samples reached full density under hot-pressing conditions.

At the maximum sintering temperatures, 1710°C and 1750°C, the SiC phase may be assumed to be non-sintering, while the Si3N 4 matrices sinter, j s Local stresses therefore develop around the SiC particles. If creep relaxation by viscous flow of the liquid phase occurs as rapidly as material densities, the densification rate is not significantly influenced by the SiC phaseJ s If creep relaxation is slower than densification, the composite densification rate becomes slower, depending upon SiC volume fraction. ~,7's With regard to the isothermal sintering of ceramic composites containing second-phase particles characterized by a much slower sintering rate than the matrix, an expression for the densification Ap/Apm.,,x was derived: Is, J9

AP/APmax= l - e x p ( 4f+gflgfl "tr) (1)

where f i f the volume fraction of the second phase, A p / A p m a x ( A p = / 9 - Pgreen; A/Omax = Pth -- Pgreen) represents the normalised fractional density at time t, Pgree, being the green density and Pth the theoretical density, r is the sintering time constant and fl is a parameter defined as the ratio of the creep rate and the densification rate during sintering. Since fl is known to be inversely proportional to the matrix viscosity, 19 both matrix grain size and liquid phase characteristics (amount and viscosity) should affect the creep capability of the Si3N 4 matrix] If the shrinkage rate of a composite is significantly reduced, increasing the volume content of the second phase, then fl becomes

much less than unity] 8'j9 If fl'> 1, the sintering behavior generally follows the rule of mixtures and the sintering time constant r remains unchanged.

Table 3 shows the time constants r and the fl values obtained from the data plotted in Fig. l(b). These estimations can be considered plausible in spite of a slight difference in the hot pressing temperatures used for monolithics and composites. Being in all cases fl,~ 1, it results that the sintering behavior of composites deviates from the rule of mixtures, j8 i.e. the SiC r~hase interferes with the mechanisms involved in the densification process. In our case, therefore, the observed densification behavior could be the result of two concurrent factors: (a) since matrix materials were pre-blended and the SiC particles subsequently added, the relative quantity of liquid phase formed in the composites during sintering was lower than in the matrices; (b) the second phase interfered, by forming a framework, with the dri- ving forces necessary for bulk diffusion around the inert particles.

Under pressureless sintering conditions, matrices and composites did not fully densify, both presenting residual porosity of some percent (Table 4). Microstructural analysis of the samples indicated strong grain growth reduction in the composites with respect to the matrices, owing to the SiC presence (Table 4). Grain growth suppres- sion in the final stage of the sintering is particularly important in order to produce ceramics with uniform microstructure. However, in this case the composites were sintered at higher temperature than matrix materials, without inducing acceler- ated grain growth, as the second phase limited the Si3N 4 grain growth by slowing down grain boundary movement through second-phase particle interaction and pore-limited migration.

Figures 2(a) and (b) and 3(a) and (b) compare the morphologies of the fracture surfaces of the hot-pressed samples, A and AC, B and BC, respectively; Figs 4(a) and (b) show fracture surfaces of pressureless sintered samples AC and

Table 3. Parameters for the study of the effect of SiC on the sintering behavior

Sample ~ (rain) fl

AC 3.36 0.053 <fl-<0.073 BC 2-84 0.057-<fl-<0.074 DC 8.52 0.046_<fl<0.078

Silicon nitride-silicon carbide composite materials

Table 4. Microstructural characteristics and mechanical properties

217

Sample Cycle p/p,~ Mean grain size a/(a + fl) E HI/~ Kt, (%) (pro) Si3N 4 (GPa) (mPa m t/e)

T(°CO t (rain)

(MPa)

R T 1000°C 1300°C

Hot-pressed A 1710 60 100"0 0.8 -- 301 2110+_92 4"80+_0-15

AC 1750 20 99"5 0'4 0"34 318 2039-+51 4"36+0'12 B 1710 60 99-7 0"6 -- 312 1825+_41 5"54+_0"29

BC 1750 20 99"5 0"2 0'28 335 2008_+61 4'46_+0"09 D 1710 60 100"0 0"4 -- 315 1977_+61 5'20+0-08

DC 1750 60 99'1 0"3 0"00 335 2069+-51 4"73+_0"15 Sintered

A 1800 120 96"0 1"9 -- -- 1254_+ 133 6-30+0"13 AC 1800 120 96"3 0-7 0"00 276 1417+_71 4'92_+0-13 B 1800 120 97"6 0-9 . . . .

BC 1800 120 97"9 0'5 0-00 296 1570+_41 5-19+_0"33

895+_35 603_+39 281_+22 786_+83 441_+34284_+22

1051-+101 683+86 571-+18 985+-41 746_+48 544_+15

1084_+48 884_+50 722_+38 756+-66 635_+12450_+13

528_+16 - -- 477+-93 -- -- 574_+29 -- -- 702+-95 -- --

BC; Figs 5(a) and (b) compare microstructures of the hot-pressed materials D and DC.

Generally, the samples obtained from the Stark Si3N 4 powder indicated a relatively wide grain- size distribution, while the UBE Si3N4 powder gave rise to a fine and uniform grain morphology, as also previously observed. 2°'21

Figure 6, a backscattered electron image of a polished surface, shows the SiC particle distribution (dark areas), Si3N4 (gray areas) and secondary phases (white areas). Most of SiC particles have submicrometric size, and very few are larger than 1 pm. Secondary phases and SiC distributions look satisfactorily uniform and regular.

The larger grain growth in the pressureless sintered materials was obviously caused by the higher temperatures and the longer soaking times which were necessary to sinter without pressure.

X-ray diffraction analysis indicated a complete a~fl-Si3N4 conversion in pressureless sintered samples, while an amount of about 30% of a- Si3N 4 was detected in the hot-pressed materials. Crystalline secondary phases were only observed in the composite DC, specifically small amounts of SizN20 and Y-A1-Si-O phases.

In any case, phase compositions, grain size and morphology were strictly related to the properties of the starting powders and to the powder treat- ment methods, as they affect the liquid phase characteristics during the sintering and the conse- quent phenomena that govern the microstructure evolution.

Mechanical properties

Young's modulus In the hot-pressed matrices, the Young's modulus (E) slightly changes from one matrix to the other

(Table 4). The lowest value was found in matrix A and the highest in matrix D with the matrix B very close to matrix D. Both the highest values were measured in the materials made with UBE powder. A clear trend of Young's modulus with some physical features, such as grain size or sintering aid amount, could not be ascertained due to the limited range of variation of these parameters. When SiC particles were added to the matrix, the Young's modulus of the composite was improved due to the higher stiffness of the reinforcing particles. Taking 410 GPa as the Young's modulus for the SiC phase, one can see that the experimental values of all the hot-pressed composites are close to the lower bound as calculated for a particulate composite. 22 The experimental values of Young's modulus of the UBE composites, i.e. BC and DC, are in good agreement with those reported by Lange. s

In the pressureless sintered composites, the Young's moduli were affected by the porosity present in these materials, as they were much lower than in the corresponding hot-pressed materials. The composite BC, being denser than AC, is slightly stiffer.

Hardness Also for this property, the experimental data show differences among the hot-pressed matrices them- selves. In this case, the highest value was measured on matrix A followed by matrices C and B, respectively. The grain size is known to have an inverse dependence upon the hardness, 23,24 but the trend in our matrices does not support this point. It could be likely that the different amount, as well as the different chemical composition of the intergranular phase, has in some way counter-


Fig. 2. Scanning electron micrograph showing the fracture surfaces of the hot-pressed matrix A (a) and composite

AC (b).

Fig. 3. Scanning electron micrograph showing the fracture surfaces of hot-pressed matrix B (a) and composite BC (b).

balanced the positive effect of reducing the grain size. The introduction of SiC particles made the composites slightly harder in most cases. Only for matrix A is the composite less hard than the starting matrix for some unexpected reason.

As for the Young's modulus, the residual porosity in the pressureless sintered materials nega- tively affected the hardness that was much lower than in the respective hot-pressed materials. In these materials, a small increase is found when SiC particles were added to the matrix (see matrix A and the relative composite AC in Table 4).

Fracture toughness The Kjc values of the hot-pressed matrices B and D are very close to each other and higher than matrix A in spite of the coarser microstructure of the latter. 25 This is probably due to the higher aspect-ratio of the grains in materials B and D with respect to the more uniform microstructure in A. Introducing SiC particles lowers the toughness in all the composites. As explained by Lange, s if the second-phase particles are below a critical size, as in our case, toughening mechanisms such as crack pinning are not operative. An ulterior factor for toughness reduction in the


Fig. 4. Scanning electron micrograph showing the fracture surfaces of the pressureless sintered composite AC (a) and

composite BC (b).

Fig. 5. Scanning electron micrograph showing the fracture surface of the hot-pressed matrix D (a) and composite DC

(b).

composites can also derive from the finer grain size of the matrix 25 when SiC particles are present (see Table 4).

For the pressureless sintered composites, the toughness values are slightly higher than the corresponding hot-pressed composites, in spite of their porosity. Considering the lower stiffness of these materials, it can be argued that porosity plays a role in the toughening mechanisms, 26 as it causes either crack wandering or crack blunting. Alternatively, since the Kxc was measured with the indentation method, it is likely that, above a

certain level of porosity, the physical mechanisms underlying the cracking process, i.e. the residual stresses induced by plastic deformation, do not hold any more. In fact, in a porous solid, part of the displaced material at the indentation site can be accommodated by densification through the crushing of pore walls (and not by plastic deformation), without any contribution to the residual stresses. The smaller the residual stresses, the smaller the extent of the crack from the corners of the indentation giving in this way unrealistically high values of fracture toughness.


Fig. 6. Backscattered electron micrograph of polished surface of the hot-pressed composite BC showing matrix grain morphology, silicon carbide particles and intergranular

amorphous phase.

owing to the low sintering aid content. Generally, the matrices were stronger than the composites with the exception of composite BC tested at 1000°C.

In all the systems investigated, the evident decrement that was found at high temperature is due to the relaxation of the residual stresses intro- duced by machining 27 as well as the generation of new defects by oxidation. 28 From this point of view, the reinforcing second phase seemed to have no influence on composite strengthening at high temperature, indicating that the mechanical behavior at high temperature was still governed by the matrix characteristics.

CONCLUSIONS

Flexural strength Room temperature. Among the hot-pressed matrices, the UBE silicon nitrides gave the highest strength results (Table 4). The finer microstructure, the high purity of the starting powders and the higher fracture toughness values compared to the STARK material could have been the factors for the improved fracture strength. The SiC addition weakened the composite in all cases. A cause for the strength reduction is, of course, the lower fracture toughness of the composite with respect to the matrix. But, as in the case of matrix D, the strength reduction, being much larger than the corresponding fracture toughness reduction, indi- cates that the second-phase particles had also changed the critical flaw size of the matrix, for example by introducing particle agglomerates that can act as fracture origins, s

In the sintered materials, almost no difference was found between the matrices. The flexural strength of these materials is much lower than the corresponding hot-pressed ones due to the residual porosity. When SiC particles were added to the matrix A, a strength reduction was observed as in the case of the hot-pressed composites, but for the matrix B a strong increment in strength was measured.

High temperature. When tested at high temperature, all the materials showed the same trend with a progressive deterioration of flexural strength as the test temperature was raised. The most refractory material was proved to be matrix D, which still had a strength of some 700 MPa at 1300°C

Si3N4-20 vol. % SiC composites, containing submicrometer-sized a-SiC particles, were produced by hot-pressing and pressureless sintering, starting from different commercial Si3N 4 powders. The effects of second phase, sintering aid amount and mixing methods on microstructure and mechanical properties were evaluated.

Under hot-pressing conditions, all the samples fully densified, and no evident densification reduction caused by the second-phase presence resulted during the heating up. The Ray and Bordia model was assumed in order to study the interaction of the SiC second phase with the mechanisms involved in the densification process. The influence of the starting powder characteristics was apparent by the different grain size distribution obtained from the two Si3N 4 powders. In any case, SiC grains had submicronic size and their distribution, together with the secondary phase distribution, was uniform and regular.

Generally, strong grain growth inhibition, in spite of the higher sintering temperature, was apparent in all the composites, owing to the SiC presence, which limited the Si3N 4 grain growth by slowing down the grain boundary movement.

The hot-pressed materials had generally better mechanical properties than the corresponding pressureless sintered ones. Among the different hot-pressed Si3N4 matrices, the most evident difference was the improvement of the high- temperature flexural strength when low amounts of sintering aids were employed. Adding 20 vol. % of SiC particles to the Si3N 4 matrices reduced, in most cases, the flexural strength of the composites both at room and high temperature, with minor


influence on the other mechanical properties investigated.

ACKNOWLEDGEMENTS

This research was carried out in the frame of the Targeted Project 'Special Materials for Advanced Technologies' of the National Research Council of Italy.

REFERENCES

1. Buljan, S. T., Baldoni, J. C. & Huckabee, M. L., Si~N4-SiC composites. Am. Ceram. Soc. Bull., 66 (2) (1987) 347-52.

2. Shalek, P. D., Petrovic, J. J., Hurley, G. E & Gac, F. D., Hot-pressed SiC-whiskers/Si3N 4 matrix composites. Am. Ceram. Soc. Bull., 65 (2)(1986) 351-6.

3. Wei, G. C. & Becher, P. E, Development of SiC whisker- reinforced ceramics. Am. Ceram. Soc. Bull., 64 (2) (1985) 298-304.

4. Bellosi, A. & De Portu, G., Hot pressed Si3N4-SiC whiskers composites. Mat. Sci. & Eng., A109 (1989) 357-62.

5. Kodoma, H., Suzuki, F., Sakamoto, H. & Miyoshi, T., Toughening of silicon nitride matrix composites by the addition of both silicon carbide whiskers and silicon carbide particles. J. Amer. Ceram. Soc., 73 (3) (1990) 678-83.

6. Akimune, Y., High strength SiC-particle/Si3N 4 composites. J. Jap. Ceram. Soc. (Int. Ed.), 98 (438) (1990) 10-13.

7. Greil, P., Petzow, G. & Tanaka, H., Sintering and HIP- ping of silicon nitride-silicon carbide composite materials. Ceram. Int., 13(1987) 19-25.

8. Lange, F. F., Effect of microstructure on strength of Si3N4-SiC composite system. J. Am. Ceram. Soc., 56 (9) (1973) 445-50.

9. Sawaguchi, A., Todha, K. & Niihada, K., Mechanical and electrical properties of silicon nitride-silicon carbide and nanocomposite materials. J. Amer. Ceram. Soc., 74 (5)(1991) 1141-4.

10. Greskovich, C. & Palm, J. A., Observations on the fracture toughness of fi-Si3N4-fl-SiC composites. J. Amer. Ceram. Soc. Disc. and Notes, Sept, Oct. (1990) 597-99.

11. Dongliang, J., Jihong, S., Shouhong, T., Jingkun, G. & Greil, P., SiC-Si3N 4 gradient composite ceramics by special hip processing. Proc. 4th Int. Syrup. on ceramic materials and components for engines, Goteborg, Sweden, ed. R. Carlsson, T. Johansson & L. Kahlman. Elsevier Applied Science, London, UK, 1992, pp. 416-23.

12. Bellosi, A., Guicciardi, S. & Tampieri, A., Development and characterization of electroconductive Si3N4-TiN composites. J. Europ. Ceram. Soc., 9 (1992) 83-93.

13. McMurtry, C. H., Boecker, W. G. D., Seshadry, S., Zanghi, J. S. & Gamier, J. E., Microstructure and material properties of SiC-TiB 2 particulate composites. Amer. Ceram. Soc. Bull., 60 (1987) 325-9.

14. Endo, H., Ueki, M. & Kubo, H., Microstructure and mechanical properties of hot-pressed SiC-TiC composites. J. Mater. Sci., 26 (1991) 3769-74.

15. ASTM C 848-78 (Reapproved 1983), Standard test method for Young's modulus and Poisson's ratio for ceramic whitewares by resonance. American Society for Testing and Materials, Philadelphia, USA.

16. Anstis, G. R., Chantikul, P., Lawn, B. R. & Marshall, D. B., A critical evaluation of indentation technique for measuring fracture toughness: I, Direct crack measurements. J. Amer. Ceram. Soc., 64 (1981) 533-8.

17. Galassi, C., Biasini, V. & Bellosi, A., Effects of powder characteristics and mixing processes on the microstructure and properties of silicon nitride. Proc. Adv. Ceram. (in press).

18. Ray, R., Bordia, R. K. & Ray, R., Sintering behaviour of bimodal powder compacts. Acta Metall., 32 (7)(1984) 1003-19.

19. Bordia, R. K. & Ray, R., Analysis of sintering of a composite with a glass or ceramic matrix. J. Amer. Ceram. Soc., 69 (3)(1986) C55-7.

20. Beilosi, A. & Galassi, C., Sintering characteristics of SigN 4 powders. High Temp. High Press., 20 (3)(1988) 335-43.

21. Biasini, V., Dalle Fabbriche, D., Graziani, T. & Bellosi, A., Influence of powder characteristics and sintering methods on silicon nitride microstructurc. Mater. Eng., 2 (1991) 373-96.

22. Ashby, M. F. & Jones, D. R. H., Engineering Materials. Pergamon Press, Oxford, UK, 1980.

23. Mukhopadhyay, A. K., Datta, S. K. & Chakraborty, D., Hardness of silicon nitride and sialon. Ceram. Int., 17 (1991) 121-7.

24. Babini, G. N., Bellosi, A. & Galassi, C., Characterization of hot-pressed silicon nitride-based materials by microhardness measurements. J. Mater. Sci., 22 (1987) 1687-93.

25. Kawashima, T., Okamoto, H., Yamamoto, H. & Kita- mura, A., Grain size dependence of the fracture toughness of silicon nitride ceramics. J. Ceram. Soc. Japan Int. edn, 99 (1991) 310-13.

26. Rice, W. R., Ceramic matrix composite toughening mechanisms: An update. Ceram. Eng. Sci. Proc., 6 (1985) 589-607.

27. Johnson-Walls, D., Evans, A. G., Marshall, D. B. & James, M. R., Residual stresses in machined ceramic surfaces. J. Amer. Ceram. Soc., 69 (1986) 44-7.

28. Evans, A. G., Engineering property requirements for high performances ceramics. Mater. Sci. & Eng., 71 (1985) 3-21.

silicon nitride-silicon carbide composite materials

Documents