《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_Physico-chemistry of interfaces in inorganic-matrix composites

COMPOSITES SCIENCE AND TECHNOLOGY ELSEVIER Composites Science and Technology 61(2001)355-362 www.elsevier.com/locate/compscitech Physico-chemistry of interfaces in inorganic-matrix composites J. Bouix *, M.P. Berthet, F. Bosselet, R. Favre, M. Peronnet, O. Rapaud, J.C. Viala C. Vincent.H.Ⅴ Incent Laboratoire des Multimateriaux et Interfaces, UMR CNRS 5615, Universite Claude Bernard Lyon 1, 43 bvd 11 Novembre 1918, F69622 villeurbanne Cedex. france Received 26 May 1999: received in revised form 20 October 1999; accepted 2 May 2000 Abstract The performances of metal matrix composites(MMCs)or ceramic matrix composites(CMCs)are mited by the char acteristics of the fibre/matrix interface or more generally those of the interfacial zone. Concerning M optimization of this zone involves control of the chemical reactivity between the reinforcement and the matrix, which te usually an out-o quilibrium system. In the case of CMCs, it is possible to obtain a non-brittle material by associating two brittle components and to exhibit a good resistance to oxidation. The physical chemist is able to offer a significant contribution for solving these problems by ting on the reinforcement surface, the matrix composition or the manufacturing conditions of the composite.@ 2001Elsevier Science Ltd. All rights reserved 如mm小oMmu 1. Introduction a dissolution-growth process and does not ave an protecting effect on the fibre which is attacked in deptl All of the recent studies show that the mechanical and Consequently, the mechanical properties fall away and thermomechanical behaviour of ceramic and metallic- the composite is highly sensitive to corrosion by humid matrix composites depends widely on the nature of the air with emission of methane. It is therefore necessary to interfacial bonding which forms between the reinforce- limit this reactivity as far as possible. Conversely ment, consisting for instance of fibres, and the matrix between the same fibres and molten magnesium, neither chemical reactivity nor wettability are noticed(Fig. 1b) Generally speaking, this bonding must be strong In this last case, it will be necessary to create a strictly nough to provide good load transfer from the matrix controlled reactivity at the fibre/matrix interface [4] to the fibres, but weak enough to deflect cracks along In ceramic-matrix composites (CMCs), chemical the interface and to avoid their propagation through the reactivity between the fibre and the matrix can also be fibre with a brittle failure of the composite an important topic. For example, when carbon or silicon In ceramic-fibre-reinforced metallic-matrix compo- carbide fibres(Nicalon)are associated with an oxide sites(MMCs), the strength of the interfacial bond base matrix, interfacial oxidation-reduction reactions depends generally on the chemical interactions occurring may proceed at high temperature. In C/C, Sic/Sic or composite. For instance, low-graphitized carbon fibres no longer take place. However, the tion reactions can between fibre and matrix during the fabrication of the C/SiC CMCs, such oxidation-reduction reactions can x-PAN T 300 or ex-Pitch P 55) have a high tensile area remains a preferential path for oxygen diffusion and strength but they are highly reactive with oxygen and therefore constitutes a weak point concerning the resis- with metals like aluminium giving the carbide Al4C3 As tance to oxidation In the case of composites working in shown in Fig. la, this carbide forms as large crystals by air and at high temperature, it is therefore important to make use of oxidation-resistant or self-repairable inter phases. Nevertheless, the main function of the interface consists in conferring a non-brittle behaviour on materials 0266-3538/01/S. see front matter C 2001 Elsevier Science Ltd. All rights reserved PII:S0266-3538(00)00107-X

Physico-chemistry of interfaces in inorganic-matrix composites J. Bouix *, M.P. Berthet, F. Bosselet, R. Favre, M. Peronnet, O. Rapaud, J.C. Viala, C. Vincent, H. Vincent Laboratoire des MultimateÂriaux et Interfaces, UMR CNRS 5615, Universite Claude Bernard Lyon 1, 43 bvd 11 Novembre 1918, F69622 Villeurbanne Cedex, France Received 26 May 1999; received in revised form 20 October 1999; accepted 2 May 2000 Abstract The performances of metal matrix composites (MMCs) or ceramic matrix composites (CMCs) are usually limited by the char￾acteristics of the ®bre/matrix interface or more generally those of the interfacial zone. Concerning MMCs, the optimization of this zone involves control of the chemical reactivity between the reinforcement and the matrix, which constitute usually an out-of￾equilibrium system. In the case of CMCs, it is possible to obtain a non-brittle material by associating two brittle components and to exhibit a good resistance to oxidation. The physical chemist is able to o€er a signi®cant contribution for solving these problems by acting on the reinforcement surface, the matrix composition or the manufacturing conditions of the composite. # 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Carbon ®bres; A. Ceramic-matrix composites (CMCs); A. Metal-matrix composites (MMCs); B. Interfaces; E. Chemical vapour deposition (CVD) 1. Introduction All of the recent studies show that the mechanical and thermomechanical behaviour of ceramic and metallic￾matrix composites depends widely on the nature of the interfacial bonding which forms between the reinforce￾ment, consisting for instance of ®bres, and the matrix [1±3]. Generally speaking, this bonding must be strong enough to provide good load transfer from the matrix to the ®bres, but weak enough to de¯ect cracks along the interface and to avoid their propagation through the ®bre with a brittle failure of the composite. In ceramic-®bre-reinforced metallic-matrix compo￾sites (MMCs), the strength of the interfacial bond depends generally on the chemical interactions occurring between ®bre and matrix during the fabrication of the composite. For instance, low-graphitized carbon ®bres (ex-PAN T 300 or ex-Pitch P 55) have a high tensile strength but they are highly reactive with oxygen and with metals like aluminium giving the carbide Al4C3. As shown in Fig. 1a, this carbide forms as large crystals by a dissolution-growth process and does not have any protecting e€ect on the ®bre which is attacked in depth. Consequently, the mechanical properties fall away and the composite is highly sensitive to corrosion by humid air with emission of methane. It is therefore necessary to limit this reactivity as far as possible. Conversely, between the same ®bres and molten magnesium, neither chemical reactivity nor wettability are noticed (Fig. 1b). In this last case, it will be necessary to create a strictly controlled reactivity at the ®bre/matrix interface [4]. In ceramic-matrix composites (CMCs), chemical reactivity between the ®bre and the matrix can also be an important topic. For example, when carbon or silicon carbide ®bres (Nicalon) are associated with an oxide base matrix, interfacial oxidation-reduction reactions may proceed at high temperature. In C/C, SiC/SiC or C/SiC CMCs, such oxidation-reduction reactions can no longer take place. However, the ®bre/matrix interfacial area remains a preferential path for oxygen di€usion and therefore constitutes a weak point concerning the resis￾tance to oxidation. In the case of composites working in air and at high temperature, it is therefore important to make use of oxidation-resistant or self-repairable inter￾phases. Nevertheless, the main function of the interface consists in conferring a non-brittle behaviour on materials 0266-3538/01/$ - see front matter # 2001 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(00)00107-X Composites Science and Technology 61 (2001) 355±362 www.elsevier.com/locate/compscitech * Corresponding author

J. Bouix et al. Composites Science and Tech 61(2001)355-3062 the fibre through thin layers deposited by CVD and working as diffusion barriers, and on the matrix com- position or the processing conditions. In CMCs, we have used single or multilayers for controlling the strength of the interfacial bonding and having a good resistance to oxidation 2. Deposit of thin refractory layers on carbon fibres by RCVD The difficulty consists in achieving a thin coating on each individual filament of a bundle constituted of several thousands of single filaments of some micrometer dia 2 um meter. The efficiency of a coating as a diffusion barrier depends on its continuity and its regular thickness along the whole length of each filament. Specific preferential deposit on the external filaments to the detriment of those situated in the tow centre must be avoide The reactive CVD (or RCVD) is seen as a promising way of achieving the surface treatment and to obtain a carbide coating(M,C: SiC, TiC, B C)on carbon fibres The coatings were prepared by heating the fibres in a gas stream carrying hydrogen and the m element of the M,C (SiCl4, TiCla or BCl3, for instance), carbon being taken on the fibre itself. For a similar kind of carbon fibre, the thickness of the coating depends much more upon the nature of the carbide and on temperature than on reaction time. Typically, the RCvd time is about 1 min. The carbide coating grows by carbon diffusion from the fibre through the layer formed already, therefore continuity Fig 1. Chemical behat a filament in two diffe and its regularity are reached even when normal pressu metallic melts: (a) strong intera ALC, formation after is used for the deposition immersion for 15 min at 680C in nium;(b)no reaction and The optimal conditions can be foreseen by thermo no wetting after 5 h immersion at 730.C in pure magnesium. dynamic calculations. The method is based on the total ibb's energy minimization of the MCl- H2/C(graphite) resulting from the coupling of two constituents which systems for a given set of conditions(temperature, gas separately exhibit brittle failure. In an axial tensile test phase composition, mole number of carbon in contact for instance, the interfacial zone must deviate in mode II with I mol of the gas mixture. )and the theoretical results the cracks induced in the matrix, thus deferring the are corroborated with experiments on bulk graphite sub- failure of the fibres and that of the composite itself. This strates and on carbon fibres with different micro- 'mechanical fuse effect can be obtained only if the structures(ex-Pan and ex-Pitch). A detailed description interfacial bonding is not too strong, which allows the of the fibre coating equipment has been given in previous activation of energy-consuming phenomena like fibre/ publications [4-6 matrix decohesion, interfacial sliding or broken fibre The uniformity and the continuity of the coating are extraction On the other hand, if the interfacial bonding confirmed by sEM observation of the oxide shells becomes too weak. a loss of contact and load transfer obtained after oxidation of as-coated carbon fibres in occurs between fibre and matrix. Optimization of the air at a temperature higher than 600oC. The residue interface requires a compromise in which the residual shown in Fig. 2 corresponds to a T300 fibre coated with thermal stress has an important part a Sic layer of 50 nm, after complete consumption of In the Laboratoire des Multimateriaux et Interfaces, carbon. The photograph indicates a thin and con- ve work to optimize the interfaces in fibre-reinforced tinuous shell that replicates the crenulated morphology metallic- and ceramic-matrix composites In MMCs, the of the fibre. This observation is proof of a continuous problem consists mainly in controlling the interfacial che- initial carbide layer. The RCV technique has been mical reactivity by acting both on the surface properties of used to fabricate more complex protective coatings such

resulting from the coupling of two constituents which separately exhibit brittle failure. In an axial tensile test for instance, the interfacial zone must deviate in mode II the cracks induced in the matrix, thus deferring the failure of the ®bres and that of the composite itself. This `mechanical fuse' e€ect can be obtained only if the interfacial bonding is not too strong, which allows the activation of energy-consuming phenomena like ®bre/ matrix decohesion, interfacial sliding or broken ®bre extraction. On the other hand, if the interfacial bonding becomes too weak, a loss of contact and load transfer occurs between ®bre and matrix. Optimization of the interface requires a compromise in which the residual thermal stress has an important part. In the Laboratoire des MultimateÂriaux et Interfaces, we work to optimize the interfaces in ®bre-reinforced metallic- and ceramic-matrix composites. In MMCs, the problem consists mainly in controlling the interfacial che￾mical reactivity by acting both on the surface properties of the ®bre through thin layers deposited by CVD and working as di€usion barriers, and on the matrix com￾position or the processing conditions. In CMCs, we have used single or multilayers for controlling the strength of the interfacial bonding and having a good resistance to oxidation. 2. Deposit of thin refractory layers on carbon ®bres by RCVD The diculty consists in achieving a thin coating on each individual ®lament of a bundle constituted of several thousands of single ®laments of some micrometer dia￾meter. The eciency of a coating as a di€usion barrier depends on its continuity and its regular thickness along the whole length of each ®lament. Speci®c preferential deposit on the external ®laments to the detriment of those situated in the tow centre must be avoided. The reactive CVD (or RCVD) is seen as a promising way of achieving the surface treatment and to obtain a carbide coating (MnC: SiC, TiC, B4C) on carbon ®bres. The coatings were prepared by heating the ®bres in a gas stream carrying hydrogen and the M element of the MnC (SiCl4, TiCl4 or BCl3, for instance), carbon being taken on the ®bre itself. For a similar kind of carbon ®bre, the thickness of the coating depends much more upon the nature of the carbide and on temperature than on reaction time. Typically, the RCVD time is about 1 min. The carbide coating grows by carbon di€usion from the ®bre through the layer formed already, therefore the coating formation is self-regulated and its continuity and its regularity are reached even when normal pressure is used for the deposition. The optimal conditions can be foreseen by thermo￾dynamic calculations. The method is based on the total Gibb's energy minimization of the MClx/H2/C(graphite) systems for a given set of conditions (temperature, gas phase composition, mole number of carbon in contact with 1 mol of the gas mixture...) and the theoretical results are corroborated with experiments on bulk graphite sub￾strates and on carbon ®bres with di€erent micro￾structures (ex-Pan and ex-Pitch). A detailed description of the ®bre coating equipment has been given in previous publications [4±6]. The uniformity and the continuity of the coating are con®rmed by SEM observation of the oxide shells obtained after oxidation of as-coated carbon ®bres in air at a temperature higher than 600C. The residue shown in Fig. 2 corresponds to a T300 ®bre coated with a SiC layer of 50 nm, after complete consumption of carbon. The photograph indicates a thin and con￾tinuous shell that replicates the crenulated morphology of the ®bre. This observation is proof of a continuous initial carbide layer. The RCVD technique has been used to fabricate more complex protective coatings such Fig. 1. Chemical behaviour of a P55 carbon ®lament in two di€erent metallic melts: (a) strong interaction with Al4C3 formation after immersion for 15 min at 680C in pure aluminium; (b) no reaction and no wetting after 5 h immersion at 730C in pure magnesium. 356 J. Bouix et al. / Composites Science and Technology 61 (2001) 355±362

J. Bouix et al. /Composites Science and Technology 61(2001)355-362 as B,C/SiC and B4C/TiB, double layers [7]. in order to fabrication of performant aluminium matrix composites increase the fibre resistance against oxidation and to with a fibre-volume fraction of 0.50 by a squeeze-casting ensure good wetting of the fibres by liquid aluminium. technique. The tensile strength is multiplied by a factor The process involves two successive RCVD steps of 3. In these materials aluminium carbide is no The presence of these thin carbide coatings is able to detected at the fibre /metal interface. w down considerably the gasification of carbon fibres The technique is not limited to modify only the carbon during an oxidation exposure and their reactivity with fibre surface, it has been applied to surface treatment of liquid aluminium. For instance, the curves of thermo- Hi-Nicalon fibres: a thin layer of Si3 N4 has been gravimetric analysis (TGA)shown in Fig 3 confirm the obtained by reaction between silicon carbide and low oxidation resistance of the pristine T300 fibre ammonia gas at a temperature higher than 1000C heated under oxygen atmosphere at 600oC. They also prove that BC single layer and B4C/SiC double layer have better protective behaviour than SiC single layer 3. Coatings with a double function gainst oxidation. Furthermore, the use of fibres coated by boron carbide or by silicon carbide has permitted The protective coatings used as interphases in MMCs re generally carbides, nitrides or oxides, i.e. brittle materials which crack at a low level of strength. when the coating thickness is up to a critical value 'ecrit it has been shown that the tensile strength of fibres decreases when the coating thickness increases. The ecrit value depends on fibre type and the adhesion strength between the fibre and the coating; however, it does not depend on the coating composition. Typically, ecrit about 16 nm for a T300 fibre. These very thin coatings are protective only when the fabrication technique of MMCs requires a very short contact time between the fibre and the heated metal When a thicker interphase is ed, the same pro- blen tered in MMCs and cmcs. i.e. the crack formation in the brittle component and the pro pagation of the cracks in the fibres. It appears that the presence of a deflector, or a 'mechanical fuse' in brittle properties of the two kinds of composites. It is possible to deflect cracks either at the nanometric scale. for instance between the graphitic planes of pyrolytic carbon Fig. 2. SEM micrograph of the residue obtained from Sic-coated or turbostratic boron nitride, or at the macroscopic T300 fibre after complete consumption of carbon by oxidation scale, for instance at different interfaces in weakly bonded multi-layered coatings 3. 1. Double layered coating: pyrocarbon/carbide We have developed a new generation of a fibre coat ing. It consists of two stacked layers: a preliminary deposition of pyrocarbon(pyC) by low-pressure CVD technique on the fibre followed by a partial conversion of this carbon layer into carbide by RCvd treatment [8]. The very thin carbon layer between the fibre and the external carbide layer acts like a mechanical fuse. Fibres with a such double-layered coating are chemically inert and are more mechanically resistant than the pristine DO(pristine) T300(SiC) fibres. Data for the tensile tests performed on three 12 fibres and four coatings can be found in Table Time(h) M40 fibres have been coated with a pyrocarbon/silicon Fig 3. Weight losses of T300 fibres coated with carbides as a function carbide(pyC/SiC) dual layer. They have been used as a of time([=600C, pO2=I bar) reinforcing agent in an aluminium matrix ID-composite

as B4C/SiC and B4C/TiB2 double layers [7], in order to increase the ®bre resistance against oxidation and to ensure good wetting of the ®bres by liquid aluminium. The process involves two successive RCVD steps. The presence of these thin carbide coatings is able to slow down considerably the gasi®cation of carbon ®bres during an oxidation exposure and their reactivity with liquid aluminium. For instance, the curves of thermo￾gravimetric analysis (TGA) shown in Fig. 3 con®rm the low oxidation resistance of the pristine T300 ®bre heated under oxygen atmosphere at 600C. They also prove that B4C single layer and B4C/SiC double layer have better protective behaviour than SiC single layer against oxidation. Furthermore, the use of ®bres coated by boron carbide or by silicon carbide has permitted fabrication of performant aluminium matrix composites with a ®bre-volume fraction of 0.50 by a squeeze-casting technique. The tensile strength is multiplied by a factor of 3. In these materials, aluminium carbide is not detected at the ®bre/metal interface. The technique is not limited to modify only the carbon ®bre surface, it has been applied to surface treatment of Hi-Nicalon ®bres: a thin layer of Si3N4 has been obtained by reaction between silicon carbide and ammonia gas at a temperature higher than 1000C. 3. Coatings with a double function The protective coatings used as interphases in MMCs are generally carbides, nitrides or oxides, i.e. brittle materials which crack at a low level of strength. When the coating thickness is up to a critical value `ecrit' it has been shown that the tensile strength of ®bres decreases when the coating thickness increases. The ecrit value depends on ®bre type and the adhesion strength between the ®bre and the coating; however, it does not depend on the coating composition. Typically, ecrit is about 16 nm for a T300 ®bre. These very thin coatings are protective only when the fabrication technique of MMCs requires a very short contact time between the ®bre and the heated metal. When a thicker interphase is required, the same pro￾blems are encountered in MMCs and CMCs, i.e. the crack formation in the brittle component and the pro￾pagation of the cracks in the ®bres. It appears that the presence of a de¯ector, or a `mechanical fuse' in brittle interphase is essential for increasing the mechanical properties of the two kinds of composites. It is possible to de¯ect cracks either at the nanometric scale, for instance between the graphitic planes of pyrolytic carbon or turbostratic boron nitride, or at the macroscopic scale, for instance at di€erent interfaces in weakly bonded multi-layered coatings. 3.1. Double layered coating: pyrocarbon/carbide We have developed a new generation of a ®bre coat￾ing. It consists of two stacked layers: a preliminary deposition of pyrocarbon (pyC) by low-pressure CVD technique on the ®bre followed by a partial conversion of this carbon layer into carbide by RCVD treatment [8]. The very thin carbon layer between the ®bre and the external carbide layer acts like a mechanical fuse. Fibres with a such double-layered coating are chemically inert and are more mechanically resistant than the pristine ®bres. Data for the tensile tests performed on three ®bres and four coatings can be found in Table 1. M40 ®bres have been coated with a pyrocarbon/silicon carbide (pyC/SiC) dual layer. They have been used as a reinforcing agent in an aluminium matrix 1D-composite Fig. 2. SEM micrograph of the residue obtained from SiC-coated T300 ®bre after complete consumption of carbon by oxidation. Fig. 3. Weight losses of T300 ®bres coated with carbides as a function of time (t=600C, pO2=1 bar). J. Bouix et al. / Composites Science and Technology 61 (2001) 355±362 357

358 J. Bouix et al. /Composites Science and Technology 61(2001)355-362 fabricated by medium-pressure foundry method. The valves opening and closing. It is possible to vary the obtained composite has a tensile strength of 1300 MPa, relative thickness of each layer by modifying the number close to the value predicted by the mixture law of pulses and the residence time of the precursors in the reactor. The seM photographs in Figs. 4 and 5 corre- 3. 2. Multi-layered coatings with different components spond, one to a coating consisted in seven(pyc/TiC) sequences and the other to a An improvement of the interphase quality consists in For these two examples of coatings, the first and the last promoting the multi-fissuration in the interphase and sub-layers are pyC. These photographs also show the deflecting the cracks parallel to the fibre; therefore the weak bonding between the different sub-layers after the idea of increasing the rupture surface number by a fracture of the fibre stacking of sequences of two layers of different comp The tensile strength value of the as-coated FT500 nents, one constituted of pyC (or hexagonal BN), the fibres(3300 MPa) is close to that of the pristine fibre other constituted of carbide. This idea has been pro-(3560 MPa)and confirms the property of mechanica posed by Diefendorf et al. [9 who have studied the effect of a coating consisted in a(py C/SiC)sequence on fibres. There exists few examples of multi-layer coatings available in literature. Further, the idea of multi-layers has been extended for the fabrication of matrices [10] Table 2 collects some data [10-15 A laminated coating constituted of distinct thin sub- layers of carbides, insulated or not by thin carbon layers is the new concept that we develop now. We can give the example of a(C/TiC)n coating on FT500 fibre. The process relies on the Pressure pulsed-CVD and the RCVD techniques. It consists in performing a series of cycles within a short period of time: each cycle includes etting the reactor under vacuum, gas injection and deposition reaction. The reactor is filled periodically with a C3H/Ar mixture or with a TiCl4/H2 one and evacuated after a chosen reaction time. The titanium carbide coating is performed by RCVD process, i. e by consumption of a part or the whole of the pyC layer which is used as the reactive substrate Gas introduction into the reactor and its evacuation are performed by pneumatic valves, and a controller is used to monitor Table I Tensile strengths (MPa) of pyC (LPCVD)/ carbide (RCVD)-coated Pristine pyc pyC/SiC pyC/TiC pyC/B4C fibres 3150 42503200 2700 34003200 560 43003900 3480 Table 2 Examples of multilayers available in literature References Interphases pyC/SiC [,12] BN/SIC [3 [14 Matrix pyC/SiC 5

fabricated by medium-pressure foundry method. The obtained composite has a tensile strength of 1300 MPa, close to the value predicted by the mixture law. 3.2. Multi-layered coatings with di€erent components An improvement of the interphase quality consists in promoting the multi-®ssuration in the interphase and de¯ecting the cracks parallel to the ®bre; therefore the idea of increasing the rupture surface number by a stacking of sequences of two layers of di€erent compo￾nents, one constituted of pyC (or hexagonal BN), the other constituted of carbide. This idea has been pro￾posed by Diefendorf et al. [9] who have studied the e€ect of a coating consisted in a (pyC/SiC) sequence on ®bres. There exists few examples of multi-layer coatings available in literature. Further, the idea of multi-layers has been extended for the fabrication of matrices [10]. Table 2 collects some data [10±15]. A laminated coating constituted of distinct thin sub￾layers of carbides, insulated or not by thin carbon layers, is the new concept that we develop now. We can give the example of a (C/TiC)n coating on FT500 ®bre. The process relies on the Pressure pulsed-CVD and the RCVD techniques. It consists in performing a series of cycles within a short period of time: each cycle includes setting the reactor under vacuum, gas injection and deposition reaction. The reactor is ®lled periodically with a C3H8/Ar mixture or with a TiCl4/H2 one and evacuated after a chosen reaction time. The titanium carbide coating is performed by RCVD process, i.e. by consumption of a part or the whole of the pyC layer which is used as the reactive substrate. Gas introduction into the reactor and its evacuation are performed by pneumatic valves, and a controller is used to monitor valves opening and closing. It is possible to vary the relative thickness of each layer by modifying the number of pulses and the residence time of the precursors in the reactor. The SEM photographs in Figs. 4 and 5 corre￾spond, one to a coating consisted in seven (pyC/TiC) sequences and the other to a (TiC/TiC)n multi-layer. For these two examples of coatings, the ®rst and the last sub-layers are pyC. These photographs also show the weak bonding between the di€erent sub-layers after the fracture of the ®bre. The tensile strength value of the as-coated FT500 ®bres (3300 MPa) is close to that of the pristine ®bre (3560 MPa) and con®rms the property of mechanical Table 1 Tensile strengths (MPa) of pyC (LPCVD)/carbide (RCVD)-coated ®bres Pristine ®bres pyC pyC/SiC pyC/TiC pyC/B4C T300-99 3150 4250 3200 2700 3800 M40 B 2740 3400 3200 3750 ± FT500 3560 4300 3900 ± 3480 Table 2 Examples of multilayers available in literature References Interphases pyC/SiC [11,12] BN/SiC [13] BN/Si3N4 [14] Matrix pyC/SiC [10,15] Fig. 5. Fractured surface of (TiC/TiC)-coated ®bre. Fig. 4. Fractured surface of (pyC/TiC)-coated ®bre. 358 J. Bouix et al. / Composites Science and Technology 61 (2001) 355±362

J. Bouix et al. /Composites Science and Technology 61(2001)355-362 of the very thin pyc layers or the stacking of when SiC is heated with pure aluminium, an invar- weakly-bonded layers of TiC. The SEM photograph in iant transformation involving four phases occurs at Fig. 6 demonstrates the presence of microcracks in the 650+3 C, i.e. at 10oC below the melting point of the carbide layers and their deflection at C/TiC interfaces metal. This transformation, which is of the quasi-peri tectic type, can be written a 3.3. Multilayers with different microstructures SiC+ Alsz2Al4C3+ L Another concept is proposed to produce fibre coat- ings and interphases in CMCs. It rests on a stacking of where Als designates metallic aluminium in the solid layers of the same chemical composition, providing that state and Lo an aluminium-rich Al/Si/ c liquid contain- successive layers have a different microstructure, for ing 1.5 at. of silicon and about I at ppm of carbon instance, a lamellar structure and an isotropic one BN is a component which fits that interphase type [16] According to the temperature during the deposition by low-pressure CVD process(LPCVD), boron nitride exhibits a more or less pronounced microtexture and hence, mechanical property anisotropic. We have shown that the use of a furnace with temperature gradients allows the deposition of a stacking of isotropic and anisotropic layers by a continuous process The SEM photographs shown in Figs. 7 and & corre spond to Hi-Nicalon fibre coated with a Bn trilayer deposited under the following conditions: BF3/NH mixture, fibre speed of 0.5 m/h. The fractured section observation of the same fibre demonstrates a decohesion between the coating and the fibre, between isotropic and anisotropic layers and in the isotropic layer The weak bonding between the fibre and coating causes the conservation of the fibre mechanical propertie 4. Interface reactivity control in carbon/aluminium and carbon/magnesium composites Fig. 6. Polished section of fibres coated by a(pyC Tic) multi-layer 4.1. Carbon/aluminium composites As pointed out in the Introduction, the main problem 1.2 in these composites is to avoid an excessive degradation of the reinforcing fibres by chemical reaction with the metal matrix during fabrication by melt-infiltration. To solve this problem, thin layers of the refractory carbides SiC, TiC and B, C have been deposited at the surface of the fibres by the rcvd process previously described. Pres- sure-infiltration of these coated fibres by liquid aluminium has resulted in composites with improved mechanical properties, showing thereby that the carbide coatings could effectively protect the underlying fibre from alu minium attack. To acquire a thorough understanding of nis protecting effect and render possible a better control of the chemical reactivity at the matrix/coating interface a detailed investigation of the chemical interactions in the al c/si Ti and Al/ c/B ternary systems has been carried out a thermodynamic approach of the chemical interac- tions in the al/c/si system under atmospheric pressure has revealed two important features [17]: Fig. 7. BN tri-layered coating

fuse of the very thin pyC layers or the stacking of weakly-bonded layers of TiC. The SEM photograph in Fig. 6 demonstrates the presence of microcracks in the carbide layers and their de¯ection at C/TiC interfaces. 3.3. Multilayers with di€erent microstructures. Another concept is proposed to produce ®bre coat￾ings and interphases in CMCs. It rests on a stacking of layers of the same chemical composition, providing that successive layers have a di€erent microstructure, for instance, a lamellar structure and an isotropic one. BN is a component which ®ts that interphase type [16]. According to the temperature during the deposition by low-pressure CVD process (LPCVD), boron nitride exhibits a more or less pronounced microtexture and, hence, mechanical property anisotropic. We have shown that the use of a furnace with temperature gradients allows the deposition of a stacking of isotropic and anisotropic layers by a continuous process. The SEM photographs shown in Figs. 7 and 8 corre￾spond to Hi-Nicalon ®bre coated with a BN trilayer deposited under the following conditions: BF3/NH3 mixture, ®bre speed of 0.5 m/h. The fractured section observation of the same ®bre demonstrates a decohesion between the coating and the ®bre, between isotropic and anisotropic layers and in the isotropic layer. The weak bonding between the ®bre and coating causes the conservation of the ®bre mechanical properties. 4. Interface reactivity control in carbon/aluminium and carbon/magnesium composites 4.1. Carbon/aluminium composites As pointed out in the Introduction, the main problem in these composites is to avoid an excessive degradation of the reinforcing ®bres by chemical reaction with the metal matrix during fabrication by melt-in®ltration. To solve this problem, thin layers of the refractory carbides SiC, TiC and B4C have been deposited at the surface of the ®bres by the RCVD process previously described. Pres￾sure-in®ltration of these coated ®bres by liquid aluminium has resulted in composites with improved mechanical properties, showing thereby that the carbide coatings could e€ectively protect the underlying ®bre from alu￾minium attack. To acquire a thorough understanding of this protecting e€ect and render possible a better control of the chemical reactivity at the matrix/coating interface, a detailed investigation of the chemical interactions in the Al/C/Si, Al/C/Ti and Al/C/B ternary systems has been carried out. A thermodynamic approach of the chemical interac￾tions in the Al/C/Si system under atmospheric pressure has revealed two important features [17]: . when SiC is heated with pure aluminium, an invar￾iant transformation involving four phases occurs at 6503C, i.e. at 10C below the melting point of the metal. This transformation, which is of the quasi-peri￾tectic type, can be written as: SiC ‡ Als ÿ!ÿAl4C3 ‡ L0 …1† where Als designates metallic aluminium in the solid state and L0 an aluminium-rich Al/Si/C liquid contain￾ing 1.5 at.% of silicon and about 1 at.ppm of carbon; Fig. 6. Polished section of ®bres coated by a (pyC /TiC) multi-layer. Fig. 7. BN tri-layered coating. J. Bouix et al. / Composites Science and Technology 61 (2001) 355±362 359

J. Bouix et al. Composites Science and Technology 61(2001)355-362 silicon to the aluminium matrix prior to infiltration such a situation will for example occur at 730@C with Al/Si alloys containing more than 4.5 at. of silicon. The phase diagram of the ternary system Al/C/Ti[18] is characterized by the existence at 812+15C of a quasi- peritectic invariant transformation which can be written on heating as: Al3Ti+Al,C3 2 Li+ TiC (3) where Li designates an aluminium base liquid containing 0.3 at. of titanium and TiCx a carbon-rich composi- tion in the homogeneity range of the titanium carbide content in the liquid is very low). As a consequence of this transformation, carbon-rich titanium carbide coating will be in thermodynamic equilibrium with aluminium base liquids at temperatures higher than or equal to 812C but will be decomposed by solid or liquid alumi nium into Al3Ti and Al4C3 at any temperature lower than 812C according to the equation 3TIC+13Al- 3Al3T1+ Al4C3 Fig 8. Fractured surface of BN-coated Hi-Nicalon fibre. As pressure infiltration of liquid aluminium is difficult e at temperatures higher than 650@C and up to at least to realize at temperatures higher than 812C, the TIC/Al 1000C, a three-phased equilibrium tends to be estab- couple has to be considered as a reactive couple in lished between SiC, Al4C3 and a Al/Si/C liquid L Under regard to composite manufacture. However, if the reac- a constant pressure, this three-phased equilibrium: tion corresponding to Eq.(4)cannot be thermo- dynamically stopped by adding titanium to the SiC+Al4C32L (2) aluminium matrix, this addition can modify the kinetics. Effectively, it has been shown that interaction is monovariant and the composition of the liquid L only between TiC and aluminium below 812C develops in depends on the temperature. At 650 C, this composition two successive stages: (i) a first stage which proceeds at a is that of Lo, i.e. 1.5 at. of silicon and about I at ppm fast rate and in which titanium produced in the convertion of carbon. Then, the silicon content of the liquid L reg- of TiC into AlC3 simply dissolves in aluminium;(i)a ularly increases with the temperature to attain 4.5 at second stage which proceeds at a much slower rate and at 730C and 13 at at 1000 C(the carbon content of in which Al3Ti crystallizes from a liquid saturated in L also increases but remains very low) titanium. Consequently, if liquid aluminium is saturated Existence of the quasi-peritectic transformation (1) in titanium prior to infiltration, the first fast-rate stage is implies that SiC is in thermodynamic equilibrium with avoided and one directly enters in the second slow-rate solid aluminium at any temperature lower than 650C. stage In this low temperature range, SiC coatings can then The Al/BC couple appears to be reactive at any protect ca on fibres very efficiently against aluminium temperature up to at least 1000 C. Decomposition of attack. At temperatures higher than 650C, Sic reacts boron carbide by aluminium below 868C gives only with pure aluminium in the solid(T 660C)to give solid Al4C3 and a liquid a carbide which is not Al4C3 as in the two preceding ISi alloy SiC coatings will then be damaged by reaction cases, but a ternary aluminium-boron carbide with the with pure aluminium. But on the one hand, this reaction chemical formula Al3 BC [19]. The reaction can be writ progresses at a slower rate than a direct C/Al interac- ten as tion owing to a smaller Gibbs free energy variation Further, interaction stops as soon as the silicon content 9Al+ 2B4C-3AIB2+2Al3BC in the al/si liquid phase has attained the proper com- position for the three-phased equilibrium(Eq 2)to be On the basis of these results, coating carbon fibres established. Decomposition of the SiC coatings can then with B,C does not a priori appear as a very interesting be completely avoided by adding the proper amount of solution for composite fabrication. In fact, B4C is

. at temperatures higher than 650C and up to at least 1000C, a three-phased equilibrium tends to be estab￾lished between SiC, Al4C3 and a Al/Si/C liquid L. Under a constant pressure, this three-phased equilibrium: SiC ‡ Al4C3 ÿ!ÿL …2† is monovariant and the composition of the liquid L only depends on the temperature. At 650C, this composition is that of L0, i.e.1.5 at.% of silicon and about 1 at.ppm of carbon. Then, the silicon content of the liquid L reg￾ularly increases with the temperature to attain 4.5 at.% at 730C and 13 at.% at 1000C (the carbon content of L also increases but remains very low). Existence of the quasi-peritectic transformation (1) implies that SiC is in thermodynamic equilibrium with solid aluminium at any temperature lower than 650C. In this low temperature range, SiC coatings can then protect carbon ®bres very eciently against aluminium attack. At temperatures higher than 650C, SiC reacts with pure aluminium in the solid (T660C) to give solid Al4C3 and a liquid Al/Si alloy. SiC coatings will then be damaged by reaction with pure aluminium. But on the one hand, this reaction progresses at a slower rate than a direct C/Al interac￾tion owing to a smaller Gibbs free energy variation. Further, interaction stops as soon as the silicon content in the Al/Si liquid phase has attained the proper com￾position for the three-phased equilibrium (Eq. 2) to be established. Decomposition of the SiC coatings can then be completely avoided by adding the proper amount of silicon to the aluminium matrix prior to in®ltration: such a situation will for example occur at 730C with Al/Si alloys containing more than 4.5 at.% of silicon. The phase diagram of the ternary system Al/C/Ti [18] is characterized by the existence at 81215C of a quasi￾peritectic invariant transformation which can be written on heating as: Al3Ti ‡ Al4C3 ÿ!ÿ L1 ‡ TiCx …3† where L1 designates an aluminium base liquid containing 0.3 at.% of titanium and TiCx a carbon-rich composi￾tion in the homogeneity range of the titanium carbide phase with x>0.9 (as in the former case, the carbon content in the liquid is very low). As a consequence of this transformation, carbon-rich titanium carbide coatings will be in thermodynamic equilibrium with aluminium base liquids at temperatures higher than or equal to 812C but will be decomposed by solid or liquid alumi￾nium into Al3Ti and Al4C3 at any temperature lower than 812C according to the equation: 3TiC ‡ 13Al ! 3Al3Ti ‡ Al4C3 …4† As pressure in®ltration of liquid aluminium is dicult to realize at temperatures higher than 812C, the TiC/Al couple has to be considered as a reactive couple in regard to composite manufacture. However, if the reac￾tion corresponding to Eq. (4) cannot be thermo￾dynamically stopped by adding titanium to the aluminium matrix, this addition can modify the reaction kinetics. E€ectively, it has been shown that interaction between TiC and aluminium below 812C develops in two successive stages: (i) a ®rst stage which proceeds at a fast rate and in which titanium produced in the convertion of TiC into Al4C3 simply dissolves in aluminium; (ii) a second stage which proceeds at a much slower rate and in which Al3Ti crystallizes from a liquid saturated in titanium. Consequently, if liquid aluminium is saturated in titanium prior to in®ltration, the ®rst fast-rate stage is avoided and one directly enters in the second slow-rate stage. The Al/B4C couple appears to be reactive at any temperature up to at least 1000C. Decomposition of boron carbide by aluminium below 868C gives only two solid compounds: the aluminium diboride AlB2 and a carbide which is not Al4C3 as in the two preceding cases, but a ternary aluminium-boron carbide with the chemical formula Al3BC [19]. The reaction can be writ￾ten as: 9Al ‡ 2B4C ! 3AlB2 ‡ 2Al3BC …5† On the basis of these results, coating carbon ®bres with B4C does not a priori appear as a very interesting solution for composite fabrication. In fact, B4C is Fig. 8. Fractured surface of BN-coated Hi-Nicalon ®bre. 360 J. Bouix et al. / Composites Science and Technology 61 (2001) 355±362

J. Bouix et al. Composites Science and Technology 61(2001)355-362 attacked at a slower rate than the disordered graphite constituting low. graphitized pan base fibres. BC coatings can then protect such fibres against aluminium attack during melt-infiltration. Another advantage to the use of B, C coatings would be that the ternary carbide Al3 BC eventually produced at the B. C/matrix interface is far less sensitive to hydrolysis than Al4C3 4.2. Carbon /magnesium composites Carbon fibres have been shown to be chemically inert in the presence of pure solid or liquid magnesium up to at least 730C [20]. Moreover, pure liquid magnesium Fig 9. The interface between M40B carbon filaments and a Mg/zr does not wet carbon(Fig. 1b). Then, contrarily to the alloy(0. 1 -).2 at y zr)after immersion for I5 min at 680.C: reactive preceding case, the main problem in the fabrication of wetting and formation of a ZrCr layer about 0.25 um in thickness. carbon/magnesium composites by melt-infiltration is to promote a controlled reaction at the metal/ fibre inter face in order to improve wetting and bonding Aluminium is the main addition element of the most commonly used magnesium foundry alloys AZ61 or AZ91(6 or 9 wt. of aluminium and I wt. of zinc). It was then logical to examine first the chemical behaviour of carbon fibres in the presence of liquid Mg/Al alloy Experiments realized at 730@C have shown that a che- mical reaction proceeds at the metal/fibre interface as soon as little amounts (less than lat. %)of aluminium are added to magnesium. The reaction product is not the expected aluminium carbide Al,C3, but a ternary carbide of magnesium and aluminium with the chemica Fig. 10. Crack propagation along the C/ZrCx interface in a M40B formula Al2MgC2 [21]. From an applied point of view, Mg/Zr composite promoting voluntarily such an interface reaction is not however, very interesting. Effectively, as Al4C3, the with the formation of a thin and continuous layer of ternary carbide Al, MgC? is formed via a dissolutio zirconium carbide ZrCx at the fibre surface(Fig. 9) precipitation mechanism in which deep dissolution pits Accordingly, this particular behaviour can be regarded are dug at the fibre surface, without any noticeable as a result of a reactive wetting process. The growth of improvement of wetting. Moreover, this ternary carbide the ZrCx layer proceeds by unidirectional solid-state is hydrolysed in humid air more rapidly than Al4C diffusion of carbon through it. As a consequence, Kir- Manganese is another carbide forming element often kendall voids are formed at the interface between the present in magnesium alloys at a level of about I at. % outer part of the carbon fibre and the surrounding ZrCx When carbon fibres are immersed in such liquid Mg-Mn layer. A weakened zone is then created in which cracks alloys, a continuous layer of manganese carbides can be deviated when a heavy stress is applied to the (Mn23 C6, Mns C2 and Mn, C3) is formed at the fibre material(Fig. 10). The rate at which the ZrCr layer surface via a crossed solid-state diffusion process. This grows is much slower than for Mg/Mn alloys. This renders growth process is more favourable to the preservation possible a precise control of the interface reactivity and of the mechanical properties of the fibres than a dis- consequently, of the strength of the bond between the solution/precipitation mechanism. However, carbides carbon fibre and the Mg/Zr alloy matrix formation proceeds at a rate too fast to be easily con trolled The most interesting results in terms of fibre/matrix 5. Conclusion interface tailoring have been obtained when zirconium was added to magnesium [22]. It has effectively been The optimization of the interface between fibre and observed that a liquid Mg-Zr alloy saturated in zirco- matrix is the essential condition for obtaining high per- nium(0.18 at. Zr at 730 C)can spread rapidly at the formance inorganic composite materials. With this carbon fibre surface and penetrate easily within the fila- objective, one can act on the surface state of the rein- ments of a fibre yarn, without severe damage to these forcement, on the matrix composition or on the processing filaments. Spreading and penetration are associated conditions

attacked at a slower rate than the disordered graphite constituting low-graphitized PAN base ®bres. B4C coatings can then protect such ®bres against aluminium attack during melt-in®ltration. Another advantage to the use of B4C coatings would be that the ternary carbide Al3BC eventually produced at the B4C/matrix interface is far less sensitive to hydrolysis than Al4C3. 4.2. Carbon/magnesium composites Carbon ®bres have been shown to be chemically inert in the presence of pure solid or liquid magnesium up to at least 730C [20]. Moreover, pure liquid magnesium does not wet carbon (Fig. 1b). Then, contrarily to the preceding case, the main problem in the fabrication of carbon/magnesium composites by melt-in®ltration is to promote a controlled reaction at the metal/®bre inter￾face in order to improve wetting and bonding. Aluminium is the main addition element of the most commonly used magnesium foundry alloys AZ61 or AZ91 (6 or 9 wt.% of aluminium and 1 wt.% of zinc). It was then logical to examine ®rst the chemical behaviour of carbon ®bres in the presence of liquid Mg/Al alloys. Experiments realized at 730C have shown that a che￾mical reaction proceeds at the metal/®bre interface as soon as little amounts (less than 1at.%) of aluminium are added to magnesium. The reaction product is not the expected aluminium carbide Al4C3, but a ternary carbide of magnesium and aluminium with the chemical formula Al2MgC2 [21]. From an applied point of view, promoting voluntarily such an interface reaction is not, however, very interesting. E€ectively, as Al4C3, the ternary carbide Al2MgC2 is formed via a dissolution/ precipitation mechanism in which deep dissolution pits are dug at the ®bre surface, without any noticeable improvement of wetting. Moreover, this ternary carbide is hydrolysed in humid air more rapidly than Al4C3. Manganese is another carbide forming element often present in magnesium alloys at a level of about 1 at.%. When carbon ®bres are immersed in such liquid Mg-Mn alloys, a continuous layer of manganese carbides (Mn23C6, Mn5C2 and Mn7C3) is formed at the ®bre surface via a crossed solid-state di€usion process. This growth process is more favourable to the preservation of the mechanical properties of the ®bres than a dis￾solution/precipitation mechanism. However, carbides formation proceeds at a rate too fast to be easily con￾trolled. The most interesting results in terms of ®bre/matrix interface tailoring have been obtained when zirconium was added to magnesium [22]. It has e€ectively been observed that a liquid Mg-Zr alloy saturated in zirco￾nium (0.18 at.% Zr at 730C) can spread rapidly at the carbon ®bre surface and penetrate easily within the ®la￾ments of a ®bre yarn, without severe damage to these ®laments. Spreading and penetration are associated with the formation of a thin and continuous layer of zirconium carbide ZrCx at the ®bre surface (Fig. 9). Accordingly, this particular behaviour can be regarded as a result of a reactive wetting process. The growth of the ZrCx layer proceeds by unidirectional solid-state di€usion of carbon through it. As a consequence, Kir￾kendall voids are formed at the interface between the outer part of the carbon ®bre and the surrounding ZrCx layer. A weakened zone is then created in which cracks can be deviated when a heavy stress is applied to the material (Fig. 10). The rate at which the ZrCx layer grows is much slower than for Mg/Mn alloys. This renders possible a precise control of the interface reactivity and, consequently, of the strength of the bond between the carbon ®bre and the Mg/Zr alloy matrix. 5. Conclusion The optimization of the interface between ®bre and matrix is the essential condition for obtaining high per￾formance inorganic composite materials. With this objective, one can act on the surface state of the rein￾forcement, on the matrix composition or on the processing conditions. Fig. 9. The interface between M40B carbon ®laments and a Mg/Zr alloy (0.15±0.2 at.% Zr) after immersion for 15 min at 680C: reactive wetting and formation of a ZrCx layer about 0.25 mm in thickness. Fig. 10. Crack propagation along the C/ZrCx interface in a M40B/ Mg/Zr composite. J. Bouix et al. / Composites Science and Technology 61 (2001) 355±362 361

J. Bouix et al. /Composites Science and Technology 61(2001)355-362 For metallic-matrix composites, the reinforcement [6 Bouix J, Viala JC, Vincent H, Vincent C, Ponthenier JL, Dazord matrix couple constitutes usually a chemically reactive J. Process for coating carbon fibre with a carbide. US patent 4 system and it is necessary to control this reactivity in the 859503,1989 manufacturing conditions of the material. This control [7 Vincent C, Piquero T, Berthet MP, Vincent H, Bouix J. Prepara- can be achieved through thermodynamics considera tion of B, C-Sic composites by RCVD from BCIy-SiClrH phase. Mater High Temp 1995: 13(1): 17-28 tion, choosing a couple initially close to the equilibrium. [8] Pastor S, Vincent C, Berthet MP, Oddou L, Vincent H, BouixJ The reaction is therefore limited and may be accom Aspects physicochimiques des procedes RCVD et LMTA de panied by a wetting effect if manufacture is performed aitement de surface permettant Infiltration spontanee des in the molten state. In other cases, it is achieved owing fibres de carbone par Aluminium. In: Favre JP, Vautrin A,edi- to kinetical effects, by directing the system to a slow and [9 Diefendorf RS, Boisvert RP Processing of polymeric precursors self-regulated reaction giving a continuous layer work posites. Mater Res soc ing as a diffusion barrier 988:120:157-62 For ceramic-matrix composites, the main problem lies [10 Heraut L, Naslain R, Quenisset JM. Procede de fabrication d'un in controlling the strength of the interfacial bonding aterial composite a matrice ceramique a tenacite amelioree. which must together provide a good load transfer and [11 Droillard C. Elaboration et caracterisation de composites a cracks deviation along the interface. The solution usually matrice SiC et a interphase sequence C/SiC. These de doctorat consists in depositing on the reinforcement surface one Sciences des Materiaux, Universite de bordeaux, 1993 or several thin layers different by their composition and [12 Heurtevent F. Materiaux multicouches nanosequences(PyC texture. The lamellar ones work as mechanical fuses thermostructuraux. These de doctorat Sciences des materiaux Concerning the deposition of thin single or multi- layers on multi-wire fibres, the CVD-like methods seem [13] Schmucker M, Schneider H, Chawla KK, Xu Zr, Ha Js.Thermal to be the most adequate, particularly reactive, pulsed degradation of fiber coatings in mullite-fiber-reinforced mullite ind low-pressure CVDs, with the objective of obtaining [14] Liu HY, Hsu SM. Fracture behavior of multilayer silicon nitride/ on each filament a continuous and regular layer also boron nitride ceramics. J Am Ceram Soc 1996: 79(9): 2452-7 able to work as protection against oxidation [15 Chen SH, Davis HO. Structural ceramic material having refra ry interface layers. US Patent 4 837 230, 1989 [16 Rebillat F. Caracterisation des interfaces et des materiaux din References terphases dans les CMCs. These de doctorat Sciences des Mate riaux. Universite de bordeaux. 1996 [ Chawla KK. In: Cahn RW, Haasen P, Kramer EJ, editors. [17 Viala JC, Fortier P, Bouix J Stable and metastable phase equili- Materials Science and Technology, vol. 13. Weinheim: VCH. 93.p.121-182. 2 Maruyama B, Barera EV, Rabenberg L. In: Everett RK and [18 Viala JC, Vincent C, Vincent H, Bouix J. Approche thermo- Arsenault RJ, editors. Metal matrix composites processing and namique de Interaction chimique entre Al et TiC. Mater Res interfaces. New York: Academic Press, 1991. p. 181-216 Bul1990:25:457-64 3 Rebillat F, Guette A, Debieuvre C, Goujard S, Naslain R Inter- [19] Viala JC, Bouix J, Gonzalez G, Esnouf C The chemical reactivity of aluminium with boron carbide B C. j Mater Sci 1997 33: 4559-73 duree de vie amelioree In: Lamon J, Baptiste D, editors. JNCIl 20 Viala JC, Fortier P, Claveyrolas G, Vincent H, Bouix J. Effect of rcachon(France): AMAC, 1998. p. 575-584 magnesium on the composition, microstructure and mechanical J Berthet MP, Bosselet F, Favre R. properties of carbon fibres. J Mater Sci 1991: 26: 4977-84 ent h. vincent C. Interface tailoring on fibr 21Bosselet F, Mentzen BF, Viala JC, Etoh MA, Bouix J Synthesis and structure of TrAl2MgC2. Eur J Solid State Inorg Chem 7(C6:191-205. 1998:35:91-9 5 Bouix J, Viala JC, Vincent H, Vincent C, Ponthenier JL, Dazord [22] Bouix J, Viala JC, Abiven H, Picouet L, Claveyrolas J Compo- J. Procede et dispositif de revetment de fibres de carbone par un site material combining a magnesium alloy containing zirconium arbure et fibres de carbone ainsi revetues. Brevet FI with a carbon reinforcement and its production process. US 1986:86:17157 Patent Demand 08 120 249. 1996

For metallic-matrix composites, the reinforcement/ matrix couple constitutes usually a chemically reactive system and it is necessary to control this reactivity in the manufacturing conditions of the material. This control can be achieved through thermodynamics considera￾tion, choosing a couple initially close to the equilibrium. The reaction is therefore limited and may be accom￾panied by a wetting e€ect if manufacture is performed in the molten state. In other cases, it is achieved owing to kinetical e€ects, by directing the system to a slow and self-regulated reaction giving a continuous layer work￾ing as a di€usion barrier. For ceramic-matrix composites, the main problem lies in controlling the strength of the interfacial bonding which must together provide a good load transfer and cracks deviation along the interface. The solution usually consists in depositing on the reinforcement surface one or several thin layers di€erent by their composition and texture. The lamellar ones work as mechanical fuses. Concerning the deposition of thin single or multi￾layers on multi-wire ®bres, the CVD-like methods seem to be the most adequate, particularly reactive, pulsed and low-pressure CVDs, with the objective of obtaining on each ®lament a continuous and regular layer also able to work as protection against oxidation. References [1] Chawla KK. In: Cahn RW, Haasen P, Kramer EJ, editors. Materials Science and Technology, vol. 13. Weinheim: VCH, 1993. p. 121±182. [2] Maruyama B, Barera EV, Rabenberg L. In: Everett RK and Arsenault RJ, editors. Metal matrix composites processing and interfaces. New York: Academic Press, 1991. p. 181±216. [3] Rebillat F, Guette A, Debieuvre C, Goujard S, Naslain R. Inter￾phase multicouche BN pour composites SiC/SiC aÁ teÂnacite et dureÂe de vie ameÂlioreÂe. In: Lamon J, Baptiste D, editors. JNC11. Arcachon (France): AMAC, 1998. p. 575±584. 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