Post-densification behaviour of pressureless sintered silicon nitride materials: Relationship between properties and microstructure

Ceramics lnternattonal 18 (1992) 119-130

Post-Densification Behaviour of Pressureless Sintered Silicon Nitride Materials: Relationship Between Properties and Microstructure L. Themelin, M. Desmaison-Brut, M. Billy Laboratolre de C6ramlques Nouvelles, Assocl~ au CNRS (URA 320), Facult8 des Sciences, Limoges, France

& J. Crampon Laboratmre de Structure et Propr16t6s de l'Etat Sohde, CNRS (URA 234), Unlverslt6 des Soences et Techmques, Vllleneuve d'Ascq, France (Received 8 February 1991, accepted 27 March 1991) Abstract: Two typical pressureless slntered SlsN4-based ceramics, using either

Y203 and Al203 as additives for the first grade (SIYAION type) or MgO and Al203 for the second (SiMgAION), were hot lsostatlcally pressed (HIPed) at 1700°C and 180 MPa. The effect of time on the post-denslficatlon, microstructure, mechanical properties and oxidation behavlour of both materials were investigated. The best performances are obtained for the first grade after a HIP treatment of just 30 mm has been apphed, while the characteristics of the second grade continually improve with time up to 2h

1 INTRODUCTION

improving density than those under an inert atmosphere of argon. A grain coarsening m a y also occur during the H I P treatment, which has a same unfavourable effect on mechanical properties. All these microstructural parameters have been regarded in this study, the aim o f which was initially to investigate the influence of duration of the postH I P treatment on the main properties of structural SiaN 4 ceramics. In order to elucidate the role of sintering aids, the authors have chosen two typical grades of pressureless sintered silicon nitride (SSN) available at present, using either Y203 and AI203 for the first grade (SiYA1ON type) or M g O and A120 a for the other (SiMgA1ON system). Density, microhardness, flexural strength and fracture toughness

H o t isostatic pressing (HIP) has been widely applied 1- s to improve the mechanical properties of relatively dense materials ( d > 94%). The beneficial effect of a p o s t - H I P treatment, indeed, is mainly due to the reduction o f pore size and o f internal defects by cracks healing. This treatment also results in reducing the glassy pockets by a regular redistribution o f the second phases along the grain boundaries. However, simultaneously, some partial decomposition of SiaN 4 and vaporization of the liquid phase m a y occur, as virtually proved by experiments carried out under high N 2 gas pressure which have been reported to be more efficient for ll9

Ceramtcs lnternanonal 0272-8842/92/$05 00 © 1992 Elsevier Science Publishers Ltd, England. Printed in Great Britain

120

L. Themehn, M. Desmaison-Brut, M. Billy, J. Crampon Table 1.

Composition SIYAION SiMgAION

Characteristics of the raw SiaN 4 materials

Relative density (%)

fl/ot

Addltwes (%wt)

Processing

5Y203 2'5AI2Oa 2" 5AI N 5MgAI204

Shp casting

98

100

CIP

97

100

+

(%)

fl

XR D analysis Y-N apatite Sl2N20 Y3AIs012 MgSiO 3 Si2N20

CIP, Cold isostatic pressing.

data at room temperature, as well as mechanical resistance and oxidation behaviour at high temperature have been analyzed and related to the microstructure. 2 EXPERIMENTAL

The two grades of pressureless sintered silicon nitride (SSN) were supplied by Ceramiques Techniques Desmarquest Cie, Trappes, France. Their characteristics are shown in Table 1. These starting materials vary by the type and amount of sintering aids, and by processing conditions as well. Their density was determined in water at r o o m temperature according to the Archimedean principle. The ~/fl ratio was specified by means of quantitative X-ray diffraction (XRD) analysis. It has been shown 6 that 50% o f A l 2 0 a used as an additive enters the SiaN 4 network of the first grade, thus leading to the composition Sis.7Al 0 300 3N7 7 for the fl phase. Transmission electron micrographs demonstrated that silicon nitride grains are surrounded at triple points and at grain boundaries by a glassy phase mainly consisting of Si, Y and Al, as shown by energy dispersive X-ray spectroscopy (EDX). In addition, there is an equivalent secondary crystalline phase in very small quantity at triple points. Regularly distributed as-grown dislocations were located on the longest fl grains (Fig. 1). As far as the SiMgAION material is concerned, theoretical calculations resulted in a similar composition for the fl phase (Sis.vgAl 0 2100.21N7 79),while the second phase is mostly amorphous with crystalline species (Si2N20 and MgSiO3) distributed from place to place. The specimens were blocks (20 x 20 x 35 m m 3) or small bars (3 x 4 x 2 2 m m 3) coated with boron nitride to prevent reaction with the graphite crucible used in the experiments. Hot isostatic compression tests were conducted with an A L S T H O M press at

1700°C and 180MPa under argon, the dwell time varying from 30 to 120 min. The grain size and morphology of hot isostatically pressed materials were examined on etched specimens in a scanning electron microscope and transmission electron microscopy was used for the study of intergranular phases. Vickers microhardness tests were performed with a diamond tip under 500g load for 10s. Flexural strength tests were carried out using the three-point bending technique at room temperature and up to 1400°C or after 24 h oxidation in air. The number of tested bars varied from six at room temperature to four at high temperature. Fracture toughness was estimated by the SENB technique where narrow notches are fractured by three-point bending. Oxidation kinetics were determined as a function of temperature, using a continuous recording thermobalance and cubic samples with 4 m m edges. 3 POST-SINTERING THROUGH HIP

CONSOLIDATION

Post-HIPing treatments have a beneficial effect on the densification of both materials despite profound differences illustrated in Fig. 2. A maximum density is obtained for Y-containing Si3N 4 when applying pressure (180 MPa) for 30 rain at 1700°C. On the contrary, the density of magnesiadoped materials gradually increases with time up to 2 h. In both cases, however, a complete densification could not be achieved, the best relative density values being 99-8 and 99.7% respectively (Fig.2). Thus, it clearly appears that an optimization of HIPing parameters, especially the dwell time, is essential. The reduction in residual porosity is obviously due to a pore closure mechanism through a rearrangement of the grains by means of the liquid phase which is formed during the HIP treatment at high temperature. As a consequence, the viscosity of

Post-densification behaviour of pressureless S S N materials

121 Density

SiYAION

3.30

SiMsAION

tl~oreticoldenaty

3.20

3.20 tl~oreticaldensuy

3.10 3.]10

1" o

Fig. 2.

0,5

!

2

(h)

ITime 2

(h)

Effect o f time on H I P post-denslficatlon of SIYAION

and SIMgAION materials.

(a)

(b)

Fig. 1. Transmission electron mlcrographs of the as-slntered (a) and post-HlPed (b) speomens viewingthe glassyphase (stars) and the secondary crystalhne phase (arrows). Scale bar: 1/~m the refractory intergranular phase must be decisive for consolidation kinetics. However, the maximum density obtained with yttria-doped materials and the decreasing values for longer dwell times (Fig. 2) suggest that decomposition and/or evaporation of Si3N 4 and/or the second phase must simultaneously

occur and cancel further densification. Effectively, a slight mass loss was then observed (0-05-0.07%) which contrasts with the absence of mass variation for materials in the system SiMgAION. Such a difference can only be related to the second phase and the conclusion is that the yttria phase is much more unstable than the magnesia one. Another factor of importance for explaining the difference in post-densification behaviour of the two materials, and especially their incomplete reduction in porosity, could be the rearrangement refusal of the elongated fl-SiaN 4 grains known as the 'bridge effect'. The microstructures of the two SiaN 4 grades are indeed quite different, as shown in Fig. 3. In the case of Y-containing materials, no essential rearrangement of the grains seems to occur through HIPmg, but a globularization by grain thickening was observed (Fig. 3(a,b)) which may cancel further improvement in densification. On the other hand, microstructural observations in the Mg system (Fig. 3(c,d)) show a better homogeneity in the distribution of the/:t-SIaN 4 needles size: their widths are sensibly the same as those measured for the first grade, but values are less dispersed. In addition, the average 'aspect ratio', which relates the needles' length to their width (R = L/d), is of the order of 7-8, and remains constant during the HIP treatment.

4 M E C H A N I C A L C H A R A C T E R I Z A T I O N AT ROOM TEMPERATURE Vickers mlcrohardness, flexural strength and fracture toughness have been tested at room temperature for each set of materials and the mean values with their standard deviations reported as a function of dwell time in Figs 4 and 5.

L. Themelin, M. Desmaison-Brut, M. Billy, J. Crampon

122

: :

.....

ii: 2 ~)

"

E~:~3 5

6i, L

',,-, ?EIO

(a)

(b)

(e) Fig. 3.

2 i:rJl I,JD15

(d)

SEM of polished or fracture surfaces of Y-containing SI3N 4 (a, b) or Mg cont. SI3N 4 (c, d) before and after 2 h HIP treatment at 1700°C, 180MPa (b,d).

G f (MPa)

Hv (Kg/mm " )

1000

KIc (MPa~/m)

m

800

2200.

--"

600

400

l

2100. 2000.

m

Hv

Of

gl¢

fly

Of

glc

Hv

Of

glc

Hv

of

Kit

200

as received

Fig. 4.

H I P 30 rain

HIP I h

HIP 2 h

Room temperature mechanical properties of Y-containing materials, evolution with time of the HIP treatment (1700°C, 180 M Pa).

Post-densificanon behaviour of pressureless SSN materials

123

(:If (MPa)

1000

nv l

(Kg/mm 2 )

Klc

(MPa,Jm) .9

l

800

m n

2200 .8

m

600

2100. l

.7

2000

400 Hv

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glc

Hv

! ¢~f

KIc

Hv

KIc

200

as

Fig. 5.

received

HIP 1 h

HIP 2 h

RT mechamcal properties of Mg-contalnmg SSN before and after HIP (1700°C, 180 MPa, 1 2 h).

4. 1 Microhardness For both materials, the Vickers microhardness drastically increases when samples have been subjected to the HIP treatment. There is indeed a statistically significant difference in both cases between the mean values obtained before and after HIPing. However, the latter values keep constant independently of the dwell time; and as they are quite similar whatever the kind of sintered silicon nitride may be (Y- or Mg-containing SSN), the conclusion is that their microhardness does not depend on the nature or composition of the second phase; it is only sensitive to residual porosity of the specimens.

4.2 Fracture strength As can be seen from Fig. 4, there is an important strength improvement of the order of 25 %, when Ycontaining Si3N 4 materials are subjected to a postsintering HIP treatment for short dwell time (30 min). This is confirmed by the variations in the SiMgAION system (Fig. 5) where the higher limit value observed for 2 h (af = 1060 MPa) is comparable w~th that obtained from the former grade in the best conditions. In all cases, the rupture had a mixed character w~th both inter- and intragranular fractures. No porosity has been observed which could have initiated the rupture. The evolution of af with time broadly follows denslfication kinetics (Fig. 2). However, this evolution must not only depend on reduction in total residual porosity, as shown by the mean value observed for a 2 h HIP treatment (Fig.4) which becomes equivalent to the initial strength of the raw material (about 800 MPa). In other words, the grain morphology induced by the HIP treatment and its

evolution with rime must play an important role on strength. Thus, in the case of a steady homogeneity of the microstructure, like that of Mg-containing SSN, the strength evolution with time follows that of densificatlon. However, the observed grain coarsening in Y-containing materials has a deleterious effect 7 which enhances the density decrease between 30 min and 2 h.

4.3 Fracture toughness This parameter is almost exclusively dependent on grain morphology and distribution regardless of the residual porosity. Indeed, fracture toughness remains roughly constant with time when the 'aspect ratio' does not change, like m the SIMgA1ON system (Fig. 5). On the other hand, decreasing K~c values, even lower than the initial one, have been observed for Y-SSN materials as grain coarsening proceeds with time. In any case, it is worth noting the higher toughness values (8-9MPax//m) obtained with these sihcon nitride ceramics which obviously result 8 from their needle-shaped microstructure (Fig. 3).

5 HIGH TEMPERATURE PROPERTIES The mechanical properties ofSi3N 4 ceramics at high temperatures are known to be governed by the nature and composition of the second phase and by their oxidation behaviour m air.

5. 1 Fracture strength The strength has been investigated in air at temperatures up to 1400°C for each kind of Si3N 4 ceramic by comparing the behavlour of a 1 h post-

124

L. Themelin, M. Desmaison-Brut, M. Billy, J. Crampon

HIPed material to that of the as-received SSN. The modulus of rupture curves plotted versus temperature (Fig. 6) have the classical shape already mentioned in the literature. 9- ~ In these experiments, the annealing or oxidation time during flexural tests is not long enough to modify the active flaws or to heal the machining defects, so that the initial microstructure will not be modified at moderate temperatures. This is the reason why the reinforcement role played on microstructure by the post-HIPed treatment keeps efficient up to 900-1000°C, as shown in Fig. 6. Above this temperature range, however, the strength gradually decreases due to a same softening of the mtergranular glassy phase, so that the trf values become similar for both as-received and post-

HIPed materials. At this stage ( T > l l 0 0 ° C ) , the fracture was found to be mainly intergranular as a result of the low viscosity of the grain-boundary phases. 5.2 Influence of oxidation on strength

Bars were exposed to air for 24h at a given temperature and the oxidized specimens fractured after cooling. The authors' results have been reported for both kinds of materials in Fig. 7, where fracture strength and relative mass gain values are plotted against temperature. As far as Y-containing materials are concerned, it appears that the strength improvement observed at room temperature for HIPed samples is no longer

Of (MPa)

1000

HIP

800 as - received

600

400

200

SiYAION

I

900

25

I 1000

I I100

Temperature (°C) I I I 1200 1300 1400

(a) [ --

Of(MPa)

800

600

400

SiMgAION

200 1

Temperature (°C) I

'

25

.¢~

I

I

900

1000

I

1100

I

1200

I

1300

I

1400

(b) Fig. 6.

Fracture strength versus temperature of as-received and HIPed SSN materials (180MPa, 1700°C, 1 h)' (a) SIYAION: (b) SIMgAION.

Post-densification behaviour of pressureless SSN materials

1000

125

of (MPa)

~

_

Am (%) me ~

HIP

1,5

800

received received

J~ 1,0

600

400 iI

200

SiYAION

/

0,5

,'

/

Temperalure (°C)

2•

800

900

1000

1100

1200

Ifm" 1300

1400

(a)

1000

~

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.~Pa)

800

dl

Am (%)

.

600

400

SiMgAION

200

/ I

Ty~vl~perature(°C) , 25

,!

I 900

I 1000

-41100

~ 1200

, 1300

, 1400

(b) Fig. 7.

Room temperature modulus of rupture of 24 h oxidized specimens subjected or not to a HIP treatment ( 180 MPa, 1700 C, l hi: (a) SJYA1ON, (b) S1MgA1ON

evtdent after annealing in air in the temperature range 800-1200cC. However, it is only from the interval 1200-1250°C when oxidation becomes effective, as shown by the mass gain evolution o f Fig. 7, that a strong reduction in strength occurs. The fact that the HIPed material keeps a slightly higher strength at elevated temperatures seems in relation to a better resistance to oxidation (see later). A similar evolution is observed in the case o f Mgcontaining Si3N 4 (Fig. 7). The slight decrease above 900°C m a y be linked to an internal oxidaUon o f the grain boundary phases in relation to ionic mobility. ~2 At l l00°C, the formation of a silica film temporarily heals the surface defects but, above 1200-1250°C, the oxidation effect becomes important by creating a new population o f strength limiting

flaws, i.e. pores, cavities, microcracks, which may become strength controlling. ~3.14

5.30xidatton kinetics in air 5.3.1 Yttrium-conta&ing SI3N 4 The mass variations per surface area resulting from oxidation tests in the temperature range 1300-1500°C have been plotted against time in Ftg. 8. It appears that the HIPed material is a little more resistant to oxidation up to 1400°C, but the kinetic curves display the same decelerated shape in both cases. Again, there was little difference in microstructure between the as-received and HIPed oxidized specimens. In both cases, there were a large number of

126

L. Themelin, M. Desmaison-Brut, M. Bdly, J. Crampon

0.4 AM

(Ks' re'z)

0.3

1500°C

0.2

iiiii

0.1

I / /

-

1,5o°

13000C

0 4

$

12

16

20

24 Time (h)

0.4 "~o

( Kg m "2)

0.3no HIP

0

.

2

1450 °C

~

(a) 1400°C

0.1

1350°C

1300 °C 0

~

i

l

4

8

a

12

l

16

l

20

i

24 Time (h)

Fig. 8. Oxidation kinetics m air of Y - c o n t a m m g SlsN ¢ specimens sub letted o1 not to a previous HIP treatment

pores through the oxide scale (Figs 9 and 10) and the silicon nitride substrate as well (Fig. 9(a)). The inner interface between nitride and oxide was always very irregular. Plate-like or needle-like yttrium silicates covered the oxide scale above 1350°C (Fig. 9(b)) and X-ray diffraction analysis showed the presence of Y2SiOs, Y2Si207 and cristobalite. The yttrium and aluminium concentration profiles, represented in Fig. 10 for a sample oxidized for 24 h at 1500°C, show that yttrium migrates from the bulk to the oxide scale. The maximum in yttrium concentration is observed near the inner interface; it corresponds to the formation of a low viscous aluminosihcate film which partially dissolves silicon nitride of the substrate. The oxidation products of this glassy SiYA1ON phase and the resulting nitrogen release and formation of bubbles occupy the major part of the scale up to the outer interface where yttrium and aluminium silicates precipitate. The oxidation curves plotted against time in the parabolic form have been reported in Fig. 11 for both kinds of specimens (HIPed and as-received SSN). As can be seen, a linearization is not applicable during the first hours of the reaction,

(b) Fig.

9.

Scanning electron mlcrographs of specimens oxl&zed for 24 h at 1300°C.

contrary to the results concerning oxidation of silicon nitride ceramics in the system SiYON mentioned so far in the literature.12,15.16 Here, in the case of a Y203-AI203 mixture used as a sintering aid, kinetics may be interpreted as follows. At the beginning of the oxidation, there is in fact the formation of a very fluid aluminosilicate film, containing additives and impurities, whose viscosity makes easier the oxygen penetration and nitrogen release. As this initial low viscous aluminosili-

Post-densihcatton behaviour of pressureless SSN materials

127

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,oi7<. I,..0. 15

10

5

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Oxide scale

Interface

;Y____

Bulk

i

15

:~

~ ~ )

200 I@

ig

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•

•

o

'

150.

100"

260 Fig. 10.

460

/

50-

YttrlUlo~x~lalnU~ItliiU~m 0cO~oCreSnlt04hi:~i!filesafter 0

cate film ts enriched by elements such as yttrium from the intergranular phase o f the nitride substrate, a crystallization in the form o f silicates Y2Si20 v and YaSiO5 may occur. The precipitation of these crystalline yttriosllicates acts as nucleation sites for the crystallization of SiO 2 as cristobalite in the oxide scale. The thickness and viscosity of the oxide layer increasing with time, due to its silica enrichment, the specimen is then progressively embedded in a protective scale which acts as a diffusion barrier. The parabolic constants Kp then deduced from the slopes of the straight lines obtained in Fig. 11 have been reported as a function of temperature in Arrhenius plots (Fig. 12). This results in different apparent activation energies according to oxidation temperatures, but values from both sides of the breaks (1350 and 1450~C for as-received and HIPed specimens, respectively) do not depend on the nature o f the material (435-850 and 440-780kJmo1-1) within the range of statistical confidence affecting the esttmation of the slopes. Both materials thus obey the same mechanism. The break m a y be related

-4'

'8

1'2 '

1;

2;

214

Fig. I 1. Oxidatton curves plotted against time m the parabolic

form at various temperatures

-3

1513 I Lu Kp

1451 I

1394 I

1340 129 I I Temperature (°C)

.

,

i

I

I

I

5.6

5.8

6

6.2

6.4

Fig. 12. Temperature dependence of the parabohc rate constants Kp Influenceof the HIP treatment prior to oxidation.

128

L. Themelin, M. Desmaison-Brut, M. Billy, J. Crampon

to the evolution of the intergranular phase composition. The YAG crystallization restrains the outward cationic migration and counterbalances viscosity. The HIP cycle, being equivalent to a thermal treatment, permits the redistribution of the intergranular phases and impurities cations, which explains the break temperature difference observed on the curves. 5.3.2 Magnesium-containing SisN 4 As could be expected from most of the studies concerned with silicon nitride sintered in the presence of magnesia, a parabolic oxidation is observed from the as-received and HIPed materials between 1350 and 1550°C (Fig. 13). The kinetic I0

/

to:

/.,ooc/ o / 15ooc (a)

8~

60

m ~ /

4~

14:o°C

20-

50

100-

150 -200 ~0 3~ Time (min)

100

80-

(nag.cm" )/ / o

60-

/

(b) Fig. 14. Micrographs of Mg-containing SiaN4 specimens after 24h oxidation in air at 1550°C.

40

20

50

100

150

200 250 300 Time (min)

Fig. 13. Oxidation kinetics of Mg-containing Si3N4 plotted in the parabolic form.

curves obtained in both cases are very similar; so too are the microstructures of the oxide layers which exhibit cracks and porosities (Fig. 14(a)) as well as a glassy phase just above the inner interface (Fig. 14(b)). X R D analysis of the oxide scale show the same presence of cristobalite and enstatite, while significant amounts of magnesium and traces of Ca and Fe have been detected by EDX analysis of the oxidized surface as early as 800°C.

Post-densi~catton behawour of pressureless SSN matertals 1579

1513

Ln Kp

1451

'

1394

' Temperature (°C)

-4

.'6 HIP -

~

~

129

additives, were cladless HIPed using argon as pressure gas (180 MPa) at 1700°C. It has been shown that post-HIP densification by pore filling is controlled by the viscosity and amount of the liquid phase. Controlling the dwell time is thus essential for optimizing the materials. Thus, the best density (near theoretical density) was obtained after a half an hour's treatment for S I 3 N 4 containing yttrium; longer times were not beneficial as the yttria phase tends to evaporate. On the other hand, density gradually increased by handling temperature and pressure for two hours in the case of Mg-containing Si3N,~.

8

4 I

5.4

Fig. 15.

I

5.6

(K'') I

5.8

N I

6

Mg-contammg $13N4 oxldaUon Temperature influence on the parabohc rate constants

The temperature dependence of the parabolic rate constants shows that the HIPed material is a little more resistant to oxtdatlon, especially at lower temperatures (Fig. 15). This results m an apparent activation energy 4 7 0 k J m o l - ~ higher than that (350 kJ mol-1) obtamed with the as-recetved SSN. As a matter of fact, these values are not far from those already obtamed for oxidation of Mgcontaining St3N ~ either by Singhal 17 (380 kJ m o l - ~), Babmt e t al. ~8 (430kJ mol-t), Desmaison-Brut and Bdly~9 (440 kJ m o l - ~) or by Veyret 2° (400 kJ m o l - 1). No break is observed on the Arrhenius diagram and this fact may be related to the absence ofintergranular phase recrystalhzation. The hterature gives a fairly good idea of the mechanism involved. 12'21 The oxidation results from two simultaneous processes which affect the second phase and the nitride grains. Due to high magnesium and impurities contents at the surface of the oxide scale and to the presence of cracks and porosities through the latter, it has been concluded that the rate controlling step is the outward diffusion of cation species through the gram boundaries of the substrate. At the mtrtde oxtde interface, a liquid phase is formed by dissolution of St3N 4 into magnesium silicate. This phase diffuses towards the external mterface and precipitates at the air contact leading to the formation of MgSiO 3 and S i O 2. 6 CONCLUSIONS Two typical pressureless sintered silicon nitrides, containing alumma wtth either Y203 or MgO as

A significant improvement of the mechanical properttes at room temperature has always been obtained after HIPmg. It essentially depends on both residual porosity (like microhardness) and homogeneity of the microstructure: a r and K~c to a lesser extent. The beneficial effect of HIPing on strength ts mamtained up to 1100-1200~C, above which the intergranular phase softening occurs. Lastly, the influence of post-HIP treatment on the oxidation behaviour has been tested. Although the reaction mechanisms remain the same for each kind of silicon nitride, a better resistance to oxidation has been pointed out at moderate temperatures up to 1400 C. Thts is due to a reduction of the glassy pockets by HIPlng and to a more regular dtstribution of the second phases at grain boundaries, which mmimtzes the outward diffusion of cationic species (Y, Mg, AI...). A C K N O W L E D G E M ENTS Thanks are due to Drs Bigay and Cales from Cbramtques Techniques Desmarquest, Trappes (France) for supplying the authors with the sintered silicon nitride samples used m this study. REFERENCES 1 YEHESKEL, O, GEFEN, Y & T A L I A N K E R , M , Hot lSOStatmg pressing of SI3N 4 w~th Y203 addmons J Mat Stteme, 19 (1984) 745 2 ZIEGLER, G & WOTTING, G , Post-treatment of preslntered SdlCOn nltnde by hot lSOStatlc pressing, hlt J Htgh Techn Ceramt~s, I (1985)31 3 THOREL, A , LAVAL, J Y & BROUSSAUD, D , High tcmperaturc mechamcal properties and mtergranular structure of stolons, J Ph.vstque,47 119861 353 4 BELLOSI, A , BIASINI, V & G U I C C I A R D I , S, Pio~ Euro Ceramt~~, Mau,~ttwht, 3 (1989) 216 5 ITURRIZA, I, CASTRO, F & FUENTES, M , Starer and starer-HIP of slhcon mtrlde ceramics with yttrm and alumina addmons, J Mat S~wnce, 24 (1989) 2047

L. Themelin, M. Desmalson-Brut, M. Billy, J. Crampon

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11 LAVAL, J. Y, DELAMARRE, C, A M A M R A , M C & BROUSSAUD, D , Influence of the mtergranular microstructure of substituted mtndes on h~gh temperature strength, J Mat. Sctence, 20 (1985) 381. 12 BILLY, M. & DESMAISON, J , H~gh temperature ox~datlon of sdlcon-based structural ceramics, High Temp. Techn, 4 (1986) 13

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