《复合材料 Composites》课程教学资源（学习资料）第五章陶瓷基复合材料 Oxide ceramic matrix/oxide fibre woven fabric composites exhibiting dissipative fracture behaviour.pdf_大学文库

t 1995 Elsevier Science Printed in Great Britain. All rig Oxide ceramic matrix oxide fibre woven fabric composites exhibiting dissipative fracture behaviour E. Mouchon and Ph. colomb ONERA-OM, BP 92, 92322 Chatillon, france Also CNRS, LASIR. 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-rein forced 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 fabric reinforced oxide matrix composite (Keywords: ceramic matrix composites: oxide interphase; fracture behaviour) INTRODUCTION vapour deposition, whereas in the second case, a carbon Ceramic matrix composites reinforced with continuous rich layer may be obtained by reaction between the fibre fibres are of considerable interest for structural applica and the matrix, which contains alkali (or alkaline earth) ons-6. However, in these composites applications are tions requiring improved mechanical properties at high limited by degradation of the carbon/boron nitride inter- temperatures in oxidizing and corrosive environments and energy production indu face and then of the Nicalon fibres, especially in an tries that require low density and damage tolerant mate rials. The advent of high-temperature components made Alternative routes have been already considered. New from ceramic matrix composites( CMCs)originates from concepts based on smart metallic coatings have been the ability of this class of materials to undergo a dissi proposed but their applications are limited in an oxidiz- pative damage tolerant behaviour, contrary to monolithic ing environment. In the field of CMCs, good mechan cal properties have previously been obtained for ceramics which exhibit a brittle fracture. This particular composites comprising oxide matrices(e. g. mullite, zirco- property of CMCs, rather unexpected for a material made nia)reinforced with SiC Nicalon fibre woven fabric that of two brittle constituents. results from the achievement during processing of a fibre-matrix interface that is suffi- were achieved through a sol-gel method. Chemically ciently good to allow load transfer from the matrix to the stable interphase layers deposited using sol-gel interface fibre pull-out. The improvement in toughness is then temperature up to 1200C (refs 9-13). Following the provided by several mechanisms, including matrix micro- same strategy, we report in this paper a preliminary study cracking, crack deflection/multiplication and finally fibre of our attempts to prepare oxide fibre woven fabric pull-out, which allows protection against the catastrophic reinforced mullite(and zirconia)matrices, exhibiting a failure of the material non-brittle fracture behaviour To date. most CMCs exhibiting good mechanical properties(e.g. high strength and high fracture tough EXPERIMENTAL ness), as well as dissipative fracture behaviour have beer obtained by associating Nicalon SiC fibres(Nippon Fibres and fabrics Carbon Co. with either non-oxide matrices such as silicon nitride or carbide or oxide glass-ceramic Three kinds of long, continuous, commercial oxide matrices, such as lithium and/or magnesium alumino fibres that belong to the aluminosilicate system were used ilicates. in the first case a carbon or a boron nitrid to realize woven fabri interface should be deposited on the fibres by chemical Nextel 440* fibre from 3M Corporation Sumica" fibre from Sumitomo Chem Co. Ltd: and To whom correspondence should be addressed Almax fibre from Mitsui Corporation COMPOSITES Volume 26 Number 3 1995 175

UTTERWORTH IN E M A N N Composites 26 (1995) 175-182 9 1995 Elsevier Science Limited Printed in Great Britain. All rights reserved 0010-4361/95:'$10.00 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 * To whom correspondence should be addressed 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 (Coton Fr~res Ets, Lyon, France). Characteristics of the fibres and fabrics are summarized in Table 1 and photographs in Figure I show the different kinds of interlaced yarns. Before achieving optimization of the processing parameters, it was necessary to determine if the thermal stability domains of the different types of oxide fibres are compatible with the sintering temperatures of our matrices and interphase precursors. Choice of compositions and sintering parameters Controlled sintering conditions are required to obtain simultaneously: (1) a full densification of the matrix and a low, open porosity of the composite; (2) no reaction between the fibre and the matrix; and (3) no excessive degradation of the fibres. Thermomechanical stability of oxide fibres in air. Figure 2 shows the evolution of the tensile strength values, measured at room temperature, for various oxide fibres after a 4 h heat treatment at different temperatures varying from 500 to 1500~ in air. Fibres investigated were Nextel 312 | amd 440 (3M Corp.), Sumica Table 1 Characteristics of bidirectional woven fabrics produced from oxide fibres Fibre ~ No. of fibres per Fibre diameter Surface weight Fabric (manufacturer) Fabrics yarn (/am) (g m -2) thickness (mm) Tensile Composition strength a (crystalline state) (MPa) Nextel 440 Taffeta b 390 -9-13 110 0.4 (3M) 8-Satin b 390 ~9-13 220 0.4 Sumica Taffeta 900 9-17 510 0.8 (Sumitomo) Almax 8-Satin 900 ~10 450 0.6 (Mitsui) 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 176 COMPOSITES Volume 26 Number 3 1995

Oxide ceramic matrix/oxide fibre composites: E. Mouchon and Ph. colomban to large weight losses, formation of porosity and growth 00 of B SiC grains. Thus oxide fibres may be preferred in N312 an oxidizing environment, especially for long-life appli N440 cations above1200°C The high temperature limit for the use of the Sumica a乏og Nextel 312 and Almax fibres in ment appears to be 1300 C, whereas the Nextel 440 fibre 51500 retains good mechanical strength after being heat-treated at 1400C for 4 h The processing temperature required unify fractory matrices, through a sol-gel route, anges between 950 C(boron-doped mullite), 1200oC (alumina) and 1300 c(zirconia )s. Furthermore, during pressure sintering, the presence of the woven 500 Nicalon fabric sheets delays the matrix densification by about 200° C at the higher ,,,L, ture range for densification of the composites should 1500 vary between 1100 and 1300C and the mullite matrix Plots of tensile strength, measured at room temperature Preparation of bidirectional oxide fibre woven fabric reinforced mullite matrix composites aluminosilicate(N312, Nextel 312: N440, Nextel 440; Su, and SiC(Nicalon, NLM 202)fibres. The dashed area corre- The method of preparation of our composites is a to the values usually given three-stage sol-gel process that has been widely described (upper, right limits)or in oxidizing (lower limit)atmos in previous references-. The first stage is the impreg nation of the fibre fabric layers by an interface gel precur or,consisting of a mixture of alkoxide precursors (Sumitomo Chem. Co. ) and Almax(Mitsui Corp. ) as diluted in an alcohol to obtain the required fluidity. The well as SiC Nicalon NLM 202 fibre(Nippon Carbon second step is the deposition of a fine and amorphous Co. ) We observe that the mechanical strength decreases powder of matrix precursor on the impregnated fibrous for all the fibres, but at different rates. This can be corre- reinforcement. Finally, hot-pressing of the stacked fabric lated to the nucleation and growth of the different stable plies in a graphite mould is achieved under controlled crystalline phases in the fibres. The alumina Almax and conditions(temperature, applied pressure, atmosphere) the aluminosilicate Nextel and Sumica fibres show differ At the beginning of our work, a simple two-step ent behaviours related to several steps of crystallization, process was used: only one impregnation was made on according to their composition. The trace observed for the fibre yarns before the hot-pressing/sintering stage and the Nextel 440 and Sumica fibres is due to the sta biliza no separate interface precursor layer was deposited. The tion of a glass-ceramic state(mullite crystals in a mullite- matrix precursor consisted of a mixture of the double like amorphous matrix), which can be maintained up to ester (BuO)2-Al-O-Si(oEt)3(ref. SiAl 084 Dynasil 1400C(refs 16 and 17). The same type of microstruc- Dynamit Nobel), as mullite and silica-rich glass precur ture will be found in our mullite matrix. The lower level sor, and B(OBu)3(ref. 10691; Alfa), as sintering activa observed for the Sumica fibre is due to its higher alumina tor. Although good densification of the matrix was content which leads to crystallization of corundum obtained the composites exhibited low mechanical prop- However. it must be underlined that the tensile matrix and the fibrous reinforcement, we decided to strength values for oxide fibres heat-treated above intercalate a zirconia interphase, which should remain 1200C are optimistic. After being thermally treated at inert during the processing and life cycles. Pure zirco- higher temperatures, a great number of fibres(typically nium propoxide(Zr(OnPr)4, ref. 88733; Fluka)was used 30-50%)are broken during handling before the tensile as interphase precursor. After a first deposit, the impreg strength measurement Figure 3b gives an illustration of nated woven fabrics were annealed in air at a temper composite hot-pressed at 1300C in nitrogen for 45 min. stabilize and promote crystallization of the zirconia The fibres appear to be polygonal: grains with a size close interphase. a second impregnation was then made to fill to 0.5 um are observable by scanning electron micro- up the voids resulting from shrinkage of the gel, before scopy(SEM). On the other hand, after a 1200 C heat the deposition of the mullite matrix precursor reatment in vacuum, the integrity of Nextel 440 fibres Mullite powder doped with boron or titanium oxide seems to be preserved, except at the matrix-fibre inter- 10 mol%) was prepared by rapid hydrolysis of the face, where a reaction occurred leading to their strong respective alkoxide, b(OBu)(ref. 10691: Alfa)or bonding(composites I and 2, Table 2). Comparison wa Ti(OEt)4(ref. 77142: Alfa), with mixtures of Al(OBu)3 made with the tensile strength values recorded on the (ref. 11140; Alfa)and Si(OMe)3 (ref. 87682; Fluka) SiC Nicalon NLM 202 fibres(mean diameter -15 um), diluted in isopropanol(alkoxides/2-propanol vol. ratio heat-treated under the same conditions. Our results agree 1). Very strong mechanical stirring was maintained with those of Prewo.. A drastic decrease of their during the hydrolysis-polycondensation stage(alkox mechanical strength is observed from 1000 C in oxidiz ides/water molar ratio =1: 40). Alcogels were then dried ing atmosphere, resulting from the formation of volatile under infra-red bulbs for a few hours. The obtained reaction products(primarily H, SiO and CO)and leading white xerogel was heat treated at 750 C in air for 10 h COMPOSITES Volume 26 Number 3 1995 177

Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban 2500 t" ~ ....... N312 9 e- N440 ............... / SiC = .......... .... / 1500 - ",:~- 1 / ~ -0 AJ 9 _ r E ,4 500 Nicalon I .... I .... J 500 1000 1500 Heat treatment temperature (~ Figure 2 Plots of tensile strength, measured at room temperature, versus 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 (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 references 9-12. 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, COMPOSITES Volume 26 Number 3 1995 177

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 Processing parameters and properties of bidirectional oxide fibre woven fabric-reinforced oxide matrix composites Fibre Strength Interface Sintering Porosity volume (MPa) Young's Composite precursor temp. Dwell Sintering (open) fraction at r.t. modulus Fibre Fabric no. (volume ratio) Matrix (~ (min) atm. (%) (%) (900~ (GPa) Nextel 440Taffeta 1 2TBB" + 3SiAl 1A1203 2SIO2-0.5B203 1250 60 vacuum 7 35 100 70 2 1TBB" + 9SiAl 1AI203-2SiO2-0.15B203 1200 60 vacuum 15 35 60 70 3 ZrW (1200) + ZrP modified LAS 1100 90 vacuum 4 15 100 (70) 40 Satin 4 1TBW + 1SiAl 3A1203-2SiO2-0.1TiO2 1200 45 N2 8 35 100 70 5 ZrP b (1050) + ZrP 3A1203 2SIO2-0.1B203 1200 30 N2 15 30 80 50 Almax Satin 6 ZrP b (1050) + ZrP 3A1203-2SIO2-0.1 B203 1200 45 N2 23 20 90 50 Sumica Taffeta 7 ZrW (1050) + ZrP 3A1203 2SIO2-0.1B203 1200 45 N2 25 25 100 (70) 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 178 COMPOSITES Volume 26 Number 3 1995

Oxide ceramic matrix/oxide fibre composites: E. Mouchon and Ph colomban 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 ity and bulk density determined by the water immersion method, The fibre volume fraction calculation is based on the theor- etical fibre density value and varies typically between 20 and 35%(Table 2) MECHANICAL BEHAVIOUR AND MICROSTRUCTURE Experimental techniques X-ray diffraction powder pattern recorded at room temperature, using a diffra ter(CuK radiation). The composite was gro an agate mortar Fibre tensile tests were performed using a monofila- ment 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 900C, after a 30 min stabilization dwell(heating rate 300C h) Typically, three rectangular samples(3 X 2X 36 mm 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 (S) using the formula e=6h8FiD2 in the elastic part of the load-defilec tion curve, where Fis 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 or by SEM (Cambidge Scan 200KV), Figures 3 and of the samples were observed using an optical microscop Results and discussion 30 Fibre-matrix reaction. Degradation of the oxide ure 5 X-ray diffraction patterns(CuKa radiation)of oxide fibre. fibres may be enhanced by the press inearth ions of reactive nosilicate matrix composite hot-pressed at 1200 C(composite 2, Table species such as boron, alkali and alkaline earth ions 2)(b)Almax satin/ZrO2 interphase/boron-doped mullite matrix which diffuse easily in aluminosilicate phases and h low melti atin/ZrOz interphase boron-doped mullite matrix ce pressed at 1200C(composite 5, Table 2): (d) Sumica taffeta/ZrO2 inter- temperature. No advantage- such as the formation o a carbon-rich sliding interphase layer resulting from a (composite 7, Table 2). O, Cristobalite;., orthorhombic mullite; *, reaction between the fibre and alkali(alkaline earth) monoclinic zirconia:V. a-alumina ions, as in the case of the SiC Nicalon fibres,4-may be expected. Comparison of the fibre sections shown in n order to remove alcohols and promote densification Figure 3 reveals the high reactivity of the oxide fibres The resulting amorphous fine powder was deposited above 1200%C. in accordance with the decrease of their the zirconium propoxide impregnated fabrics using ng mechanical strength(Figure 2). The fibres in a compos C6HsCI as solvent. Poly(methyl methacrylate)(PMMA) ite sintered at 1150oC maintain their integrity whereas powder (ref. 18. 224.9; Aldrich) was added to stabilize those in a composite sintered at 1300C show drastic performed under vacuum, with a nitrogen atmosphere scopic grains and corrugated surfaces are visibe micro- and adjust the viscosity of the slip. Pressure sintering was changes: the section becomes polygonal with being sometimes imposed above 400oc Table 2) When the matrix precursor comprised the mixture of Heating rate was 250 C h up to 600C and increased the double ester ( BuO)x-Al-o-Si(OEt)3(silica-rich to 350 C h-lup to the dwell temperature. a pressure of mullite glass precursor) and B(OBu)3 as sintering acti- 2.5 MPa was applied at the beginning of the hot-press- vator, good densification of the matrix was obtained but ing cycle to ensure good contact between the fibre yarns the materials exhibited low toughness(composites I and and the interphase and matrix precursors. The applied 2, Table 2). After a heat treatment at 1200-1250C in ressure was then gradually raised between 600 and vacuum, it seems that a reaction occurred between the 900 C, up to 20 MPa. This maximum value was reached two boron-rich aluminosilicate components leading to at the beginning of the mullite intrinsic shrinkage, which the formation of a low-melting point aluminoborosili is related to the dehydroxylation-nucleation reaction 8. cate interphase. The fibre-matrix interface was strong 180 COMPOSITES Volume 26 Number 3 1995

Oxide ceramic matrix~oxide fibre composites: E. Mouchon and Ph. Colomban 9 9 9 9 9 9 9 oV r 9 9 I t t 50 30 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; 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). MECHANICAL BEHAVIOUR AND MICROSTRUCTURE 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/D 2 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. Results and discussion 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, 180 COMPOSITES Volume 26 Number 3 1995

Oxide ceramic matrix/oxide fibre composites: E. Mouchon and Ph Colomban To prevent reaction between the matrix and the composites exhibiting a non-brittle fracture behaviour fibrous reinforcement, we intercalated an interphase The absence of reaction between the matrix and the fib layer between the mullite matrix and the fibrous rein- is a very important criterion, and zirconia appears to be forcement. Its composition was selected from those a good interphase material to be associated with oxide refractories that should remain inert at working temper- fibres belonging to the Al2O3-SiO2 system. Although our atures: in this paper, a zirconia interphase was consid- composites with zirconia interphase exhibited a dissipa- ered. Pure zirconium propoxide was used tive fracture behaviour, their mechanical properties precursor. The Nextel 440 satin reinforced modified remain poor(O3p=100 MPa) due to the high residual lithium aluminosilicate(LAS)matrix composite with a open porosity (15-20%)and the low fibre volume frac- zirconia interphase layer, hot-pressed at 1100 C tion(20-35%). In particular, owing to the low tempera- (composite 3, Table 2) had a low open porosi ture of processing, the sol-gel zirconia layer remains and no visible reaction appeared at the ZrO 2 lay porous or contains cracks resulting from the high shrink interface(Figures 4e, f) despite the high reactivity of this age associated with the gel-ceramic transformation. It matrix. X-ray diffraction patterns indicated that the would be interesting to know if this remaining porosity zirconia interphase has monoclinic symmetry for promotes or inhibits the dissipative character of the frac composites hot-pressed at 1100 and 1200oC(Figures 5b, ure. The mechanical properties are, however, very similar to those previously obtained for SiC Nicalon This is not surprising as it usually does not appear for woven fabric-reinforced mullite matrix composites heat treatment below 1300 C (ref. 11) before optimization of the composition of the interfacial precursors and processing parameters Work and surface fracture. Figure 6 shows repr sentative load-strain curves recorded for several oxide ACKNOWLEDGEmENTs fibre-reinforced aluminosilicate matrices, with or without intercalated zirconia interphase. When the interfa Useful discussions with Dr M. Parlier are gratefully precursor consisted of a mixture of the double ester acknowledged ( BuO)2-Al-O-Si(OEt)3, glass precursor and PO(OBu) as sintering activator, good densification of the compos- REFERENCES ite was obtained (composite 4, Table 2). Whereas neither Cooper, R F, and Chyung, KJ.. Mater. Sci. 1987, 22, 3148 in the corresponding fracture micrograph(Figure 3a). On neny, J, Van Valza, J.R. and Kelly, M.A. J. Am. Ceram. Soc. the other hand, a significant work of fracture is observe 1990,73,2054 5 Bleay, S, Scott, V.D., Harris, B, Cooke, R G. and Habib, FA on the corresponding flexural bending load-strain curves for composites with a zirconia interphase. A departure 6 Naslain.R, in ' Proc. emes Journees Nationales sur les of the linear straight line is clearly observed in Figures Composites(JNC-8),(Eds O. Allix, J F. Favre and P Ladeveze), 66, c. Multiplication of microcracks in the zirconia inter phase around the fibres, before failure of the composite, 7 Kerans, R.J. in 'High-Temperature Ceramic Matrix Composites, Proc. 6th EACM-HT-CMC Conf ' (Eds r. Naslain, J. Lamon and may explain such behaviour. a similar behaviour was D. Doumeingts), Woodhead Publishing Ltd, Cambridge, 1993, observed with the Nextel 440 satin/ZrO2 interphase/ 01-312 boron-doped mullite matrix composite(composite 5 8 Davis. J and Evans. A G. in 'H Table 2). When the fibrous reinforcement consists of es, Proc. 6th EACM-HT-CMC Conf ' (Eds R. Naslai J. Lamon and D. Doumeingts), Woodhead Publishing Ltd Almax woven yarns(composite 6, Table 2), a progres sive stress transfer between the zirconia interphase layer 9 Colomban, Ph. in Proc. &emes Journees Nationales sur and the fibre is observed on the recorded flexural bending Composites(JNC-8)(Eds O. Allix, J F. Favre and P. Ladeveze), plots(Figures 6c, d) and simultaneously fibre pull-out is shown by optical microscopy observation( Figures 7a, arlier,M. Patents FR2672283,EP92400235-5.US7.830904, b). Macroscopically, photographs of the fracture reveal that the material behaves as a lamellar composite: crack 11 Colomban, Ph and Bouchon, E. in'High-Tempe Matrix Com deflection at the planar(mullite matrix-ZrO2 interphase) posites, Proc. 6th EACM Naslain, J. Lamon and D, Doumeingts), Woo and(ZrO2 interphase-oxide fibre) interfaces was Ltd, Cambridge, 1993, pp. 159-166 observed. A gradual decrease of the tensile strength(o3p) 12 Mouchon, E. Ph. D, Thesis, University of Paris vI, 993 value with the progressive matrix/interphase-fibre stress 13 Bruneton, E, Michel, D. and Colomban, Ph. Physique 1 transfer is only well observed for the Sumica taffeta/Zro c.Bl.1989,68,39 interphase/mullite matrix composite(composite 7, Table 15 Prewo, K M. Am. Ceram. Soc. Bull. 1989. 68. 401 2), according to the photographs of the non-brittle frac- 16 Colomban, Ph and Mazerolles, L. J. Mater. Sci. Lett. 1990,9, ture(Figures 7c, d). In this case, the fibre pull-out length reaches 200 um. Simultaneously, the strain value at the 17 Low, I.M. and MacPherson, R. J. Mater. Sci. 1989, 23, 189 18 Colomban, Ph and Vendange, VJ non-Cryst. Sol. 1922, 147& rupture point reaches 0.4%(Figure 6d) 9 vendange V, and Colomban, Ph J. Sol-Gel Sci. Technol. 1994. 2,407 CONCLUSION 20 Bruneton, E and Colomban, Ph J non-Cryst. Sol. 1992, 147& 148.201 Many conditions needed to be fulfilled to obtain contin- 21 Mouchon, E and Colomban, Ph J. Physique IV(Paris)1993 oxide fibre-reinforced ceramic oxide matrix 7,1941 182 COMPOSITES Volume 26 Number 3 1995

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 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. acknowledged. Parlier are gratefully 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, EP 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 IV (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 IV (Paris) 1993, C7, 1941 182 COMPOSITES Volume 26 Number 3 1995

《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料 Oxide ceramic matrix/oxide fibre woven fabric composites exhibiting dissipative fracture behaviour

《复合材料 Composites》课程教学资源（学习资料）第五章陶瓷基复合材料 Oxide ceramic matrix/oxide fibre woven fabric composites exhibiting dissipative fracture behaviour