oxide fibre woven fabric composites exhibiting dissipative fracture behaviour

Composites 26 (1995) 175-182 9 1995 Elsevier Science Limited Printed in Great Britain. All rights reserved 0010-4361/95:'$10.00

U T T E R W O R T H IN E M A N N

Oxide ceramic matrix/oxide fibre woven fabric composites exhibiting dissipative fracture behaviour E. Mouchon and Ph. Colomban* ONERA-OM, BP 92, 92322 Chatillon, France. A/so CNRS, LAS/R, 2 rue Henri Dunant, 94320 Thiais, France (Received 28 March 1994; revised 10 June 1994)

The effects of a zirconia interphase on the mechanical properties of long, continuous, oxide fibre-reinforced oxide ceramic matrix (mullite) composites prepared through a sol-gel route have been studied by scanning electron microscopy, optical microscopy, X-ray diffraction analysis and three-point bending tests. The achievement of a non-brittle composite was related to the absence of chemical reaction between the zirconia interphase and the oxide fibres. Comparison was made with SiC Nicalon fibre woven fabricreinforced oxide matrix composites. (Keywords: ceramic matrix composites; oxide fibres; interphase; fracture behaviour)

INTRODUCTION Ceramic matrix composites reinforced with continuous fibres are of considerable interest for structural applications requiring improved mechanical properties at high temperatures in oxidizing and corrosive environments, especially in the aerospace and energy production industries that require low density and damage tolerant materials. The advent of high-temperature components made from ceramic matrix composites (CMCs) originates from the ability of this class of materials to undergo a dissipative damage tolerant behaviour, contrary to monolithic ceramics which exhibit a brittle fracture. This particular property of CMCs, rather unexpected for a material made of two brittle constituents, results from the achievement during processing of a fibre-matrix interface that is sufficiently good to allow load transfer from the matrix to the fibre, but sufficiently weak to allow a fibrous fracture with fibre pull-out. The improvement in toughness is then provided by several mechanisms, including matrix microcracking, crack deflection/multiplication and finally fibre pull-out, which allows protection against the catastrophic failure of the material 12. To date, most CMCs exhibiting good mechanical properties (e.g. high strength and high fracture toughness), as well as dissipative fracture behaviour, have been obtained by associating Nicalon SiC fibres (Nippon Carbon Co.) with either non-oxide matrices such as silicon nitride or carbide, or oxide glass-ceramic matrices, such as lithium and/or magnesium aluminosilicates. In the first case a carbon or a boron nitride interface should be deposited on the fibres by chemical

* T o w h o m c o r r e s p o n d e n c e s h o u l d be a d d r e s s e d

vapour deposition, whereas in the second case, a carbonrich layer may be obtained by reaction between the fibre and the matrix, which contains alkali (or alkaline earth) ions I-6. However, in these composites applications are limited by degradation of the carbon/boron nitride interface and then of the Nicalon fibres, especially in an oxidizing atmosphere. Alternative routes have been already considered. New concepts based on smart metallic coatings have been proposed but their applications are limited in an oxidizing environment 78. In the field of CMCs, good mechanical properties have previously been obtained for composites comprising oxide matrices (e.g. mullite, zirconia) reinforced with SiC Nicalon fibre woven fabric that were achieved through a sol-gel method. Chemically stable interphase layers deposited using sol-gel interface precursors led to good mechanical behaviour from room temperature up to 1200~ (refs 9-13). Following the same strategy, we report in this paper a preliminary study of our attempts to prepare oxide fibre woven fabricreinforced mullite (and zirconia) matrices, exhibiting a non-brittle fracture behaviour. EXPERIMENTAL

Fibres and fabrics Three kinds of long, continuous, commercial oxide fibres that belong to the aluminosilicate system were used to realize woven fabrics: 9 Nextel 440~ fibre from 3M Corporation; 9 Sumica :~ fibre from Sumitomo Chem. Co. Ltd; and 9 Almax ~ fibre from Mitsui Corporation.

COMPOSITES Volume 26 Number 3 1995 175

Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban Fibre yarns were woven into taffeta or satin fabrics ( C o t o n Fr~res Ets, L y o n , France). Characteristics o f the fibres and fabrics are summarized in Table 1 and p h o t o g r a p h s in Figure I show the different kinds o f interlaced yarns.

Before achieving optimization o f the processing parameters, it was necessary to determine if the thermal stability domains o f the different types o f oxide fibres are compatible with the sintering temperatures o f o u r matrices and interphase precursors.

Choice of compositions and sintering parameters

Thermomechanical stability of oxide fibres in air. Figure 2 shows the evolution o f the tensile strength

Controlled sintering conditions are required to obtain simultaneously: (1) a full densification o f the matrix and a low, open porosity o f the composite; (2) no reaction between the fibre and the matrix; and (3) no excessive degradation o f the fibres.

values, measured at r o o m temperature, for various oxide fibres after a 4 h heat treatment at different temperatures varying f r o m 500 to 1500~ in air. Fibres investigated were Nextel 312 | a m d 440 (3M Corp.), Sumica

Table 1 Characteristics of bidirectional woven fabrics produced from oxide fibres Fibre ~ (manufacturer)

Fabrics

No. of fibres per yarn

Fibre diameter (/am)

Surface weight (g m-2)

Fabric thickness(mm)

Nextel 440 (3M)

Taffetab 8-Satinb

390 390

-9-13 ~9-13

110 220

0.4 0.4

Sumica (Sumitomo)

Taffeta

900

9-17

510

0.8

Almax (Mitsui)

8-Satin

900

~10

450

0.6

Composition (crystalline state)

Tensile strength a (MPa)

7A1203-2.8SIO2-0.2B203 ~2000 (mullite amorph. + orthorhombic + TA1203) 8.5A1203-1.5SIO2 ~ 1800 (mullite amorph. + orthorhombic + 3/A1203) A1203 + SiO~ traces -1200 (~/A1203+ SiO2 amorph.)

a Measured at room temperature with a gauge length equal to 10 mm b Woven by Coton Fr~res Ets (Lyon, France)

Figure 1 Photographs (magnification • 0.88) of bidirectional woven fabrics of Nextel 440 (a, taffeta; c, satin), Sumica (b, taffeta) and Almax (d, satin) oxide fibres

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COMPOSITES Volume 26 Number 3 1995

Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban

2500 t " ~ ....... N312 9e - N440

=

...............

. . . . . . . . . .

1500

~

SiC

/

. . . .

",:~-

-

-0

9

/

1/

AJ

_

r

E ,4

500

Nicalon I

500

.

.

.

.

I

.

.

.

.

1000

J

1500

Heat t r e a t m e n t t e m p e r a t u r e (~ Plots of tensile strength, measured at room temperature, v e r s u s temperature of a 4 h heat treatment in air, for alumina (A1, Almax), atuminosilicate (N312, Nextel 312; N440, Nextel 440; Su, Sumiea) and SiC (Nicalon, NLM 202) fibres. The dashed area corresponds to the values usually given for SiC Nicalon fibres heated in neutral (upper, right limits)or in oxidizing (lower limit) atmospheres x4,~s

Figure 2

(Sumitomo Chem. Co.) and Almax (Mitsui Corp.), as well as SiC Nicalon NLM 202| fibre (Nippon Carbon Co.). We observe that the mechanical strength decreases for all the fibres, but at different rates. This can be correlated to the nucleation and growth of the different stable crystalline phases in the fibres. The alumina Almax and the aluminosilicate Nextel and Sumica fibres show different behaviours related to several steps of crystallization, according to their composition. The trace observed for the Nextel 440 and Sumica fibres is due to the stabilization of a glass-ceramic state (mullite crystals in a mullitelike amorphous matrix), which can be maintained up to 1400~ (refs 16 and 17). The same type of microstructure will be found in our mullite matrix. The lower level observed for the Sumica fibre is due to its higher alumina content which leads to crystallization of corundum above 1200~ However, it must be underlined that the tensile strength values for oxide fibres heat-treated above 1200~ are optimistic. After being thermally treated at higher temperatures, a great number of fibres (typically 30-50%) are broken during handling before the tensile strength measurement. Figure 3b gives an illustration of the crystallization that occurs in Almax fibres inside a composite hot-pressed at 1300~ in nitrogen for 45 rain. The fibres appear to be polygonal: grains with a size close to 0.5 txm are observable by scanning electron microscopy (SEM). On the other hand, after a 1200~ heat treatment in vacuum, the integrity of Nextel 440 fibres seems to be preserved, except at the matrix-fibre interface, where a reaction occurred leading to their strong bonding (composites 1 and 2, Table 2). Comparison was made with the tensile strength values recorded on the SiC Nicalon NLM 202 fibres (mean diameter ~15 txm), heat-treated under the same conditions. Our results agree with those of Prewo 14'1~. A drastic decrease of their mechanical strength is observed from 1000~ in oxidizing atmosphere, resulting from the formation of volatile reaction products (primarily H, SiO and CO) and leading

to large weight losses, formation of porosity and growth of/3 SiC grains. Thus oxide fibres may be preferred in an oxidizing environment, especially for long-life applications above 1200~ The high temperature limit for the use of the Sumica, Nextel 312 and Almax fibres in an oxidizing environment appears to be 1300~ whereas the Nextel 440 fibre retains good mechanical strength after being heat-treated at 1400~ for 4 h. The processing temperature required to densify refractory matrices, through a sol-gel route, ranges between 950~ (boron-doped mullite), 1200~ (alumina) and 1300~ (zirconia) 182~ Furthermore, during pressure sintering, the presence of the woven fabric sheets delays the matrix densification by about 200~ at the higher temperatures 9. Thus, the temperature range for densification of the composites should vary between 1100 and 1300~ and the mullite matrix may be preferred.

Preparation of bidirectional oxide fibre woven fabricreinforced mullite matrix composites The method of preparation of our composites is a three-stage so,gel process that has been widely described in previous references9-12. The first stage is the impregnation of the fibre fabric layers by an interface gel precursor, consisting of a mixture of alkoxide precursors diluted in an alcohol to obtain the required fluidity. The second step is the deposition of a fine and amorphous powder of matrix precursor on the impregnated fibrous reinforcement. Finally, hot-pressing of the stacked fabric plies in a graphite mould is achieved under controlled conditions (temperature, applied pressure, atmosphere). At the beginning of our work, a simple two-step process was used: only one impregnation was made on the fibre yarns before the hot-pressing/sintering stage and no separate interface precursor layer was deposited. The matrix precursor consisted of a mixture of the double ester (BuO)2-A10-Si(OEt)3 (ref. SiAl 084 Dynasil; Dynamit Nobel), as mullite and silica-rich glass precursor, and B(OBu)3 (ref. 10691; Alfa), as sintering activator. Although good densification of the matrix was obtained, the composites exhibited low mechanical properties (Table 2). To prevent strong bonding between the matrix and the fibrous reinforcement, we decided to intercalate a zirconia interphase, which should remain inert during the processing and life cycles. Pure zirconium propoxide (Zr(OnPr)4, ref. 88733; Fluka) was used as interphase precursor. After a first deposit, the impregnated woven fabrics were annealed in air at a temperature between 1050 and 1200~ in order to densify, stabilize and promote crystallization of the zirconia interphase. A second impregnation was then made to fill up the voids resulting from shrinkage of the gel, before the deposition of the mullite matrix precursor. Mullite powder doped with boron or titanium oxide (I0 tool%) was prepared by rapid hydrolysis of the respective alkoxide, B(OBu)3 (ref. 10691; Alfa) or Ti(OEt)4 (ref. 77142; Alfa), with mixtures of AI(OBu)3 (ref. 11140; Alfa) and Si(OMe)3 (ref. 87682; Fluka) diluted in isopropanol (alkoxides/2-propanol vol. ratio -- 1:1). Very strong mechanical stirring was maintained during the hydrolysis-polycondensation stage (alkoxides/water molar ratio -- 1:40). Alcogels were then dried under infra-red bulbs for a few hours. The obtained white xerogel was heat treated at 750~ in air for 10 h,

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Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban

Figure 3 SEM micrographs of: (a) the fracture of a Nextel 440 satin/interface precursor (TBP/SiAt)/titanium-doped mullite matrix composite (composite 4, Table 2) hot-pressed at 1200~ and annealed for 2 h in air at the same temperature; (b) the polished section of an Almax satin/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1300~ for 45 min; (c) the fracture of a Nextel 440 satin/mullite matrix composite hot-pressed at 1150~ for 20 min with boron oxide sintering activator; and (d) the polished section of a Sumica taffeta/ZrO2 interphase/borondoped mullite composite hot-pressed at 1300~ for 45 min

Table 2

Fibre

Processing parameters and properties of bidirectional oxide fibre woven fabric-reinforced oxide matrix composites Interface Composite precursor Fabric no. (volume ratio)

Nextel 440Taffeta

Satin Almax Sumica

Satin Taffeta

1 2 3

2TBB" + 3SiAl 1TBB" + 9SiAl ZrW (1200) + ZrP

Matrix

Sintering temp. (~

Dwell (min)

1A12032SIO2-0.5B203 1AI203-2SiO2-0.15B203 modified LAS

1250 1200 1100

60 60 90

Porosity Sintering (open) atm. (%) vacuum vacuum vacuum

7 15 4

Fibre volume fraction (%)

Strength (MPa) Young's at r.t. modulus (900~ (GPa)

35 35 15

100 60 100 (70)

70 70 40 70

4

1TBW + 1SiAl

3A1203-2SiO2-0.1TiO2

1200

45

N2

8

35

100

5

ZrP b (1050) + ZrP

3A1203 2SIO2-0.1B203

1200

30

N2

15

30

80

50

6 7

ZrP b (1050) + ZrP ZrW (1050) + ZrP

3A1203-2SIO2-0.1 B203 3A1203 2SIO2-0.1B203

1200 1200

45 45

N2 N2

23 25

20 25

90 100 (70)

50 50

aAcetone is added hThe fabrics are impregnated with zirconium propoxide; after gelling, impregnated fabrics are heated for 2 h in air, at the temperature given in parentheses, before a second impregnation TBB: tributylborate, B(OBu)3 (ref. 10691: Alfa-Ventron) SiAl: (BuO)2-A1-O-Si(OEt)3 (ref. SiAl 084 Dynasil; H(ils France (formerly Dynamit Nobel)) ZrP: zirconium propoxide, Zr(OnPr)4 (ref. 88733; Fluka) TBP: tributylphosphate, PO(OBu)3 (Alfa-Ventron) LAS: 0.5Li20-A1203-0.5P2Os-0.1ZrO2-0.1TiO2

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Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban

Figure 4 SEM micrographs of bidirectional oxide fibre woven fabric-reinforced oxide matrix composites: (a) section of an Almax satin/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1200~ (composite 5, Table 2); (b) detail of (a); (c) sliced section of a Sumica taffeta/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1200~ (composite 7, Table 2); (d) detail of (c); (e) polished section of a Nextel 440 taffeta/ZrOe interphase/modified LAS matrix composite hot-pressed at 1100~ (composite 3, Table 2); (f') detail of (e)

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Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban

9

9

9

9

9

the pressure was removed just before the end of the dwell. Cooling was complete after 5 h. The composites were then roughly polished, and their open porosity and bulk density determined by the water immersion method. The fibre volume fraction calculation is based on the theoretical fibre density value and varies typically between 20 and 35% (Table 2).

9

MECHANICAL BEHAVIOUR AND MICROSTRUCTURE

9

oV

r

9

I 50

Experimental techniques X-ray diffraction powder patterns were recorded at room temperature, using a diffractometer (CuK~ radiation). The composite was ground in an agate mortar. Fibre tensile tests were performed using a monofilament tensile apparatus with a 10 mm gauge length. For each heat treatment, 30 fibres of each kind were measured and the average value is reported in Figure 2. Composite flexural strength and elastic modulus were determined by three-point bending tests. Tests were performed in air at room temperature and at 900~ after a 30 min stabilization dwell (heating rate 300~ h 1). Typically, three rectangular samples (3 • 2 x 36 mm 3) were broken for each composite at both temperatures. The strain rate was 0.3 mm min -~. The strain (e) was calculated from the bending deflection (6) using the formula e = 6h6F/D2 in the elastic part of the load-deflection curve, where F is the applied load, D is the distance between the knifes and h is the thickness of the sample. Fracture surfaces as well as sliced or polished sections of the samples were observed using an optical microscope or by SEM (Cambidge Scan 200KV), Figures 3 and 4.

9

t

t 30

Results and discussion 20

Figure 5 X-ray diffraction patterns (CuK~ radiation) of oxide fibrereinforced oxide matrix composites: (a) Nextel 440 taffeta/boroaluminosilicate matrix composite hot-pressed at 1200~ (composite 2, Table 2); (b) Almax satin/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1200~ (composite 6, Table 2); (c) Nextel 440 satin/ZrO2 interphase boron-doped mullite matrix composite hotpressed at 1200~ (composite 5, Table 2); (d) Sumica taffeta/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1200~ (composite 7, Table 2). 0, Cristobalite; O, orthorhombic mullite; "L-, monoclinic zirconia; V, a-alumina

in order to remove alcohols and promote densification. The resulting amorphous fine powder was deposited on the zirconium propoxide impregnated fabrics using C6H5C1 as solvent. Poly(methyl methacrylate) (PMMA) powder (ref. 18.224.9; Aldrich) was added to stabilize and adjust the viscosity of the slip. Pressure sintering was performed under vacuum, with a nitrogen atmosphere being sometimes imposed above 400~ (Table 2). Heating rate was 250~ h -1 up to 600~ and increased to 350~ h -l up to the dwell temperature. A pressure of 2.5 MPa was applied at the beginning of the hot-pressing cycle to ensure good contact between the fibre yams and the interphase and matrix precursors. The applied pressure was then gradually raised between 600 and 900~ up to 20 MPa. This maximum value was reached at the beginning of the mullite intrinsic shrinkage, which is related to the dehydroxylation-nucleation reaction 18,19;

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COMPOSITES Volume 26 Number 3 1995

Fibre-matrix reaction. Degradation of the oxide fibres may be enhanced by the presence of reactive species such as boron, alkali and alkaline earth ions which diffuse easily in aluminosilicate phases and promote the formation of phases with low melting temperature. No advantage - such as the formation of a carbon-rich sliding interphase layer resulting from a reaction between the fibre and alkali (alkaline earth) ions, as in the case of the SiC Nicalon fibres 1-6'21 - may be expected. Comparison of the fibre sections shown in Figure 3 reveals the high reactivity of the oxide fibres above 1200~ in accordance with the decrease of their mechanical strength (Figure 2). The fibres in a composite sintered at 1150~ maintain their integrity whereas those in a composite sintered at 1300~ show drastic changes: the section becomes polygonal with microscopic grains and corrugated surfaces are visible. When the matrix precursor comprised the mixture of the double ester (BuO)z-A1-O-Si(OEt)3 (silica-rich mullite glass precursor) and B(OBu)3 as sintering activator, good densification of the matrix was obtained but the materials exhibited low toughness (composites 1 and 2, Table 2). After a heat treatment at 1200-1250~ in vacuum, it seems that a reaction occurred between the two boron-rich aluminosilicate components leading to the formation of a low-melting point aluminoborosilicate interphase. The fibre-matrix interface was strong,

Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban a C

d

~0[ , i i , [ i0~ O, 1

(

0.1

e(%)

ol

03

~-

0.1

0.3

0.5

Figure 6 Flexural bending load versus strain curves, recorded at room temperature, in air, for bidirectional oxide fibre woven fabric-reinforced oxide matrix composites: (a) Nextel 440 taffeta/boroaluminosilicate matrix composite (1200~ (composite 2, Table 2); (b) Nextel 440 satirgZrO2 interphase/modified LAS matrix composite (1100~ (composite 3, Table 2); (c) Almax satin/ZrO~ interphase/boron-doped mullite matrix composite (1200~ (composite 6, Table 2); (d) Sumica taffeta/ZrO2 interphase/boron-doped mullite matrix composite (1200~ (composite 7, Table 2). The dashed trace corresponds to a measurement performed at 900~ in air, after a 30 rain stabilization dwell. Flexural strength and elastic modulus values are given in Table 2

Figure 7 Optical micrographs of fractures at room temperature (three-point bending tests) of: (a) Almax satin/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1200~ (composite 6, Table 2); (b) detail of (a); (c), (d) Sumica taffeta/ZrO2 interphase/boron-doped mullite matrix composite hot-pressed at 1200~ (composite 7, Table 2), top and lateral views, respectively. The scale is given by the fibre diameter (-15/xm)

giving rise to a brittle fracture behaviour. Furthermore, Nextel 400 fibres were broken and many cracks are observable by SEM (Figure 3c).This may be correlated

with the destabilization of the aluminosilicate phase and crystallization of the cristobalite phase, as shown in the corresponding X-ray spectrum (Figure 5a).

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Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban To prevent reaction between the matrix and the fibrous reinforcement, we intercalated an interphase layer between the mullite matrix and the fibrous reinforcement. Its composition was selected from those refractories that should remain inert at working temperatures: in this paper, a zirconia interphase was considered. Pure zirconium propoxide was used as interphase precursor. The Nextel 440 satin reinforced modified lithium aluminosilicate (LAS) matrix composite with a zirconia interphase layer, hot-pressed at ll00~ (composite 3, Table 2) had a low open porosity (~4% and no visible reaction appeared at the ZrO2 layer-fibre interface (Figures 4e, f) despite the high reactivity of this matrix. X-ray diffraction patterns indicated that the zirconia interphase has monoclinic symmetry for composites hot-pressed at 1100 and 1200~ (Figures 5b, c, d). No trace of zircon (ZrSiO4) phase was detected. This is not surprising as it usually does not appear for heat treatment below 1300~ (ref. 11).

Work and surface fracture. Figure 6 shows representative load-strain curves recorded for several oxide fibre-reinforced aluminosilicate matrices, with or without intercalated zirconia interphase. When the interface precursor consisted of a mixture of the double ester (BuO)z-A1-O-Si(OEt)3, glass precursor and PO(OBu)3 as sintering activator, good densification of the composite was obtained (composite 4, Table 2).Whereas neither matrix microcracking nor the fibre pull-out effect is noticed on the stress-strain curve (Figure 6a), a small length of fibre pull-out, typically 5-50 txm, was observed in the corresponding fracture micrograph (Figure 3a). On the other hand, a significant work of fracture is observed on the corresponding flexural bending load-strain curves for composites with a zirconia interphase. A departure of the linear straight line is clearly observed in Figures 6b, c. Multiplication of microcracks in the zirconia interphase around the fibres, before failure of the composite, may explain such behaviour. A similar behaviour was observed with the Nextel 440 satin/ZrO2 interphase/ boron-doped mullite matrix composite (composite 5, Table 2). When the fibrous reinforcement consists of Almax woven yarns (composite 6, Table 2), a progressive stress transfer between the zirconia interphase layer and the fibre is observed on the recorded flexural bending plots (Figures 6c, d) and simultaneously fibre pull-out is shown by optical microscopy observation (Figures 7a, b). Macroscopically, photographs of the fracture reveal that the material behaves as a lamellar composite: crack deflection at the planar (mullite matrix-ZrO2 interphase) and (ZrO2 interphase-oxide fibre) interfaces was observed. A gradual decrease of the tensile strength (Cr3p) value with the progressive matrix/interphase-fibre stress transfer is only well observed for the Sumica taffeta/ZrO2 interphase/mullite matrix composite (composite 7, Table 2), according to the photographs of the non-brittle fracture (Figures' 7c, d). In this case, the fibre pull-out length reaches 200 Ixm. Simultaneously, the strain value at the rupture point reaches 0.4% (Figure 6d). CONCLUSION Many conditions needed to be fulfilled to obtain continuous oxide fibre-reinforced ceramic oxide matrix

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COMPOSITES Volume 26 Number 3 1995

composites exhibiting a non-brittle fracture behaviour. The absence of reaction between the matrix and the fibre is a very important criterion, and zirconia appears to be a good interphase material to be associated with oxide fibres belonging to the A1203-SIO2 system. Although our composites with zirconia interphase exhibited a dissipative fracture behaviour, their mechanical properties remain poor (O-3p= 100 MPa) due to the high residual open porosity (15-20%) and the low fibre volume fraction (20-35%). In particular, owing to the low temperature of processing, the sol-gel zirconia layer remains porous or contains cracks resulting from the high shrinkage associated with the gel-ceramic transformation. It would be interesting to know if this remaining porosity promotes or inhibits the dissipative character of the fracture. The mechanical properties are, however, very similar to those previously obtained for SiC Nicalon woven fabric-reinforced mullite matrix composites before optimization of the composition of the interfacial precursors and processing parameters ~~ ACKNOWLEDGEMENTS Useful discussions with Dr M. Parlier are gratefully acknowledged. REFERENCES 1 Mah, T.I., Mendiratta, M.G., Katz, A.P. and Mazdiyasni, K.S. Am. Ceram. Soc. Bull. 1987, 66(2), 304 2 Evans, A.G.J. Am. Ceram. Soe. 1990, 73, 187 3 Cooper, R.F. and Chyung, K.J.J. Mater. Sci. 1987, 22, 3148 4 Homeny, J., Van Valza, J.R. and Kelly, M.A.J. Am. Ceram. Soc. 1990, 73, 2054 5 Bleay, S., Scott, V.D., Harris, B., Cooke, R.G. and Habib, F.A. J. Mater. Sci. 1992, 27, 281 6 Naslain, R. in 'Proc. 86rues Journ~es Nationales sur les Composites (JNC-8)' (Eds O. Allix, J.F. Favre and P. Ladev~ze), AMAC, Paris, 1992, pp. 199-221 7 Kerans, R.J. in 'High-Temperature Ceramic Matrix Composites, Proc. 6th EACM-HT-CMC ConE' (Eds R. Naslain, J. Lamon and D. Doumeingts), Woodhead Publishing Ltd, Cambridge, 1993, pp. 301-312 8 Davis, J. and Evans, A.G. in 'High-Temperature Ceramic Matrix Composites, Proc. 6th EACM-HT-CMC Conf.' (Eds R. Naslain, J. Lamon and D. Doumeingts), Woodhead Publishing Ltd, Cambridge, 1993, pp. 329-336 9 Colomban, Ph. in 'Proc. 86mes Journ6es Nationales sur les Composites (JNC-8)' (Eds O. Allix, J.F. Favre and P. Ladev6ze), AMAC, Paris, 1992, pp. 73-84 10 Colomban, Ph., Menet, M., Mouchon, E., Courtemanche, C. and Parlier, M. Patents FR 2672283, E P 92400235-5, US 7.830904, ONERA, February 1991 I1 Colomban, Ph. and Mouchon, E. in 'High-Temperature Ceramic Matrix Composites', Proc. 6th EACM-HT-CMC Conf.' (Eds R. Naslain, J. Lamon and D. Doumeingts), Woodhead Publishing Ltd, Cambridge, 1993, pp. 159-166 12 Mouchon, E. Ph.D. Thesis, University of Paris VI, 1993 13 Bruneton, E., Michel, D. and Colomban, Ph. J. Physique I V (Paris) 1993, C7, 1977 14 Prewo, K.M. Am. Ceram. Soe. Bull. 1989, 68, 395 15 Prewo, K.M. Am. Ceram. Soe. Bull. 1989, 68, 401 16 Colomban, Ph. and Mazerolles, L. J. Mater. Sci. Lett. 1990, 9, 1077 17 Low, I.M. and MacPherson, R. Z Mater. Sci. 1989, 23, 1895 18 Colomban, Ph. and Vendange, V. J. non-Cryst. Sol. 1922, 147 & 148, 245 19 Vendange, V. and Colomban, Ph. J. Sol-Gel Sci. Technol. 1994, 2, 407 20 Bruneton, E. and Colomban, Ph. J. non-Cryst. Sol. 1992, 147 & 148, 201 21 Mouchon, E. and Colomban, Ph. J. Physique I V (Paris) 1993, C7, 1941