and manufacturing ELSEVIER Composites: Part A 30(1999)537-547 Micro/minicomposites: a useful approach to the design and development of non-oxide cmcs Roger Naslain", Jacques Lamon, Rene Pailler, Xavier Bourrat, Alain Guette, Francis Langlais aboratory for Thermostructural Composites, UMR-47(CNRS-SEP-UB1), University of bordeaux. 3 Allee de la boetie, 33600 Pessac, france Micro(one single filament)and mini (one single fiber tow) non-oxide composites(C/C; C/SiC and SiC/SiC)with simple(Pyc or BN)or complex interphases [C (B)or(Pyc-SiC) multilayers] are fabricated in a short time by CVD/CVI. The fiber/matrix interfacial zone is characterized by aEs and TEM. Tensile tests are used to assess the mechanical properties and the weibull statistical parameters of both the fiber and matrix, as well as the fiber-matrix interfacial parameters(Ti Id, Gis). The tensile stress-strain behaviour has been modelled. The tensile curves exhibit the same features as those previously reported for real nD-composites. Lifetime at high temperatures in air is characterized through static/cyclic fatigue tests and modelled. It is improved by replacing conventional pyrocarbon by highly engineered interphases. The micro/mini composite approach is used in the optimization of processing conditions and to derive parameters necessary fo the modelling of the thermomechanical and chemical behaviour of composites with more complex fiber architectures. o 1999 Elsevier Science ltd. All rights reserved Keywords: A Ceramic matrix composites(CMCs); Model composites; B. Interface/interphase 1. Introduction CMC, which requires several processing/characterization be tir Non-oxide ceramic matrix composites(CMCs), such as fabrication technique. Moreover, the complexity of real C/C, C/SiC or SiC/SiC composites, usually exhibit a fiber architectures often precludes the derivation of simple complex fiber architecture(2D, 2.5D or 3D). They are correlations between composite properties and processing produced according to liquid or gas phase routes requiring conditions relatively long processing times. In the liquid phase routes, within the scope of the design and development of new the starting fibrous material is impregnated with a liquid materials it can be more appropriate to use ID model precursor of the matrix, e.g. a slurry or an organic/organo- composites, such as the microcomposites or the minicom metallic polymer, and pyrolysed at high temperature. The posites(comprising one single fiber or one single tow, impregnation/pyrolysis sequence is repeated several times respectively), in order to conduct several processing/char- in order to achieve a high densification level. In the ga acterization iterative loops in a relatively short time phase routes, such as the isothermal/isobaric chemical Furthermore, for such very simple fiber architectures vapor infiltration process(I-CVI), a porous fiber preform micromechanics-based models exist which can be used to is infiltrated with the matrix deposited from a gaseous derive useful material parameters, such as load transfer precursor(a hydrocarbon for carbon and a chlorosilane for parameters, from simple mechanical tests [5,6]. Examples SiC). The deposition process should be conducted at low of studies conducted via the use of the micro/mini compo- temperature and low pressure, in order to avoid an early site approach, have been already reported in the field of non- sealing of the pore entrances. The densification duration oxide materials, however it is not by far a generalised way can be relatively long, depending on the size of the preform for designing CMCs and for optimising their processing and the residual porosity [1]. Although accelerated CVI- conditions [7-191 processes have been proposed, the densification duration Micro/mini model composites have been used during is still of several tens of hours [2-4]. Thus, the optimizatio almost one decade at LCts to optimize the fiber-matrix of the composition and the processing conditions for a given ( FM)interfacial zone in SiC/SiC composites and to generate micromechanical data necessary for modelling the mechan- en四+3006 ical behaviour. More recently, the approach has been tended to cc co 1359-835X/99/S- see front matter @1999 Elsevier Science Ltd. All rights reserved P:S1359-835X(98)00147-X
Micro/minicomposites: a useful approach to the design and development of non-oxide CMCs Roger Naslain*, Jacques Lamon, Rene´ Pailler, Xavier Bourrat, Alain Guette, Francis Langlais Laboratory for Thermostructural Composites, UMR-47 (CNRS-SEP-UB1), University of Bordeaux, 3 Alle´e de La Boe¨tie, 33600 Pessac, France Abstract Micro (one single filament) and mini (one single fiber tow) non-oxide composites (C/C; C/SiC and SiC/SiC) with simple (PyC or BN) or complex interphases [C (B) or (PyC-SiC)n multilayers] are fabricated in a short time by CVD/CVI. The fiber/matrix interfacial zone is characterized by AES and TEM. Tensile tests are used to assess the mechanical properties and the Weibull statistical parameters of both the fiber and matrix, as well as the fiber–matrix interfacial parameters (ti; ld; Gic). The tensile stress–strain behaviour has been modelled. The tensile curves exhibit the same features as those previously reported for real nD-composites. Lifetime at high temperatures in air is characterized through static/cyclic fatigue tests and modelled. It is improved by replacing conventional pyrocarbon by highly engineered interphases. The micro/mini composite approach is used in the optimization of processing conditions and to derive parameters necessary for the modelling of the thermomechanical and chemical behaviour of composites with more complex fiber architectures. q 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic matrix composites (CMCs); Model composites; B. Interface/interphase 1. Introduction Non-oxide ceramic matrix composites (CMCs), such as C/C, C/SiC or SiC/SiC composites, usually exhibit a complex fiber architecture (2D, 2.5D or 3D). They are produced according to liquid or gas phase routes requiring relatively long processing times. In the liquid phase routes, the starting fibrous material is impregnated with a liquid precursor of the matrix, e.g. a slurry or an organic/organometallic polymer, and pyrolysed at high temperature. The impregnation/pyrolysis sequence is repeated several times in order to achieve a high densification level. In the gas phase routes, such as the isothermal/isobaric chemical vapor infiltration process (I-CVI), a porous fiber preform is infiltrated with the matrix deposited from a gaseous precursor (a hydrocarbon for carbon and a chlorosilane for SiC). The deposition process should be conducted at low temperature and low pressure, in order to avoid an early sealing of the pore entrances. The densification duration can be relatively long, depending on the size of the preform and the residual porosity [1]. Although accelerated CVIprocesses have been proposed, the densification duration is still of several tens of hours [2–4]. Thus, the optimization of the composition and the processing conditions for a given CMC, which requires several processing/characterization loops, can be time consuming whatever the nature of the fabrication technique. Moreover, the complexity of real fiber architectures often precludes the derivation of simple correlations between composite properties and processing conditions. Within the scope of the design and development of new materials it can be more appropriate to use 1D model composites, such as the microcomposites or the minicomposites (comprising one single fiber or one single tow, respectively), in order to conduct several processing/characterization iterative loops in a relatively short time. Furthermore, for such very simple fiber architectures, micromechanics-based models exist which can be used to derive useful material parameters, such as load transfer parameters, from simple mechanical tests [5,6]. Examples of studies conducted via the use of the micro/mini composite approach, have been already reported in the field of nonoxide materials, however it is not by far a generalised way for designing CMCs and for optimising their processing conditions [7–19]. Micro/mini model composites have been used during almost one decade at LCTS to optimize the fiber–matrix (FM) interfacial zone in SiC/SiC composites and to generate micromechanical data necessary for modelling the mechanical behaviour. More recently, the approach has been extended to C/C composites [20]. The aim of the present Composites: Part A 30 (1999) 537–547 1359-835X/99/$ - see front matter q 1999 Elsevier Science Ltd. All rights reserved. PII: S1359-835X(98)00147-X * Corresponding author. Tel.: 133-5-56844706; fax: 133-5-56841225; e-mail: admin@lcts.u-bordeaux.fr
R. Naslain et al/Composites: Part A 30(1999)537-547 Fig 1. Nicalon/SiC microcomposites:(a)samp ing CVD-processing,(b)morphology of a SiC/BN/SiC microcomposite failed under tensile loading at room temperature showing fiber pull-out. contribution is to show how such model composites can be prepared and characterized, on the one hand, and to give the limits of the approach, on the other hand 2. Experimental SiC/SiC micro-or minl-composites comprise one straight single filament or a tow, respectively. Most experiments were performed with Si-C-O ex-polycarbosilane Nicalon fibers or Si-C oxygen-free Hi-Nicalon fibers, both display ing 500 filaments per yarn (with a filament diameter of 14 um)and being manufactured by Nippon Carbon. Some experiments were also conducted with high strength ex- Pan carbon fibers (filament diameter of 7 um). Micro 4≥s and minicomposites were fabricated according to I-CVI or pressure pulsed CVI (P-CVI)techniques, which have been described in detail elsewhere [1, 21]. In the preparation of minicomposite, the starting material was a length of the as- received multifilament yarn, slightly twisted and maintained in straight configuration with a ceramic holder. In that of a microcomposite, a length of one single filament was care- acted from the yarn and mounted on a ceramic holder with a high temperature cement, as shown in Fi 1. In a second step, the interphase was deposited onto the fiber surface from suitable gaseous precursors including: (i) hydrocarbons, such as CH4, C3 H8, C3H, for anisotropic pyrocarbon interphases; ( ii) BF3-NH3 or BCly-NH3-H2 mixtures for the deposition of boron nitride; (iii) hydrocar- bons and CHa SiCl, /H for multilayered(Pyc-SiC)n inter- phases and (iv)C3Hg-BCl3-H2 mixtures for composition graded C(B)interphases. The interphase, with an overall 186R thickness of a few 100 nm, was usually deposited(micro- composites)or infiltrated (minicomposites) by P-CVD/P- CVI rather than by conventional I-CVI, inasmuch as the Fig. 2. C/C minicomposites:(a) failure surface; (b)multiple matrix micro- former allows a better control of the morphology and texture acking, according to Ref. [24] of the interphase. As a matter of fact, complex multilayered
contribution is to show how such model composites can be prepared and characterized, on the one hand, and to give the limits of the approach, on the other hand. 2. Experimental SiC/SiC micro- or mini-composites comprise one straight single filament or a tow, respectively. Most experiments were performed with Si-C-O ex-polycarbosilane Nicalon fibers or Si-C oxygen-free Hi-Nicalon fibers, both displaying 500 filaments per yarn (with a filament diameter of 14 mm) and being manufactured by Nippon Carbon. Some experiments were also conducted with high strength exPAN carbon fibers (filament diameter of 7 mm). Micro and minicomposites were fabricated according to I-CVI or pressure pulsed CVI (P-CVI) techniques, which have been described in detail elsewhere [1,21]. In the preparation of a minicomposite, the starting material was a length of the asreceived multifilament yarn, slightly twisted and maintained in straight configuration with a ceramic holder. In that of a microcomposite, a length of one single filament was carefully extracted from the yarn and mounted on a ceramic holder with a high temperature cement, as shown in Fig. 1. In a second step, the interphase was deposited onto the fiber surface from suitable gaseous precursors including: (i) hydrocarbons, such as CH4; C3H8; C3H6, for anisotropic pyrocarbon interphases; (ii) BF3-NH3 or BCl3-NH3-H2 mixtures for the deposition of boron nitride; (iii) hydrocarbons and CH3SiCl3/H2 for multilayered (PyC-SiC)n interphases and (iv) C3H8-BCl3-H2 mixtures for composition graded C (B) interphases. The interphase, with an overall thickness of a few 100 nm, was usually deposited (microcomposites) or infiltrated (minicomposites) by P-CVD/PCVI rather than by conventional I-CVI, inasmuch as the former allows a better control of the morphology and texture of the interphase. As a matter of fact, complex multilayered 538 R. Naslain et al. / Composites: Part A 30 (1999) 537–547 Fig. 1. Nicalon/SiC microcomposites: (a) sample holder used during CVD-processing, (b) morphology of a SiC/BN/SiC microcomposite failed under tensile loading at room temperature showing fiber pull-out. Fig. 2. C/C minicomposites: (a) failure surface; (b) multiple matrix microcracking, according to Ref. [24]
R. Naslain et al/Composites: Part A 30(1999)537-547 fixed epoXy mini SiC CVD calibrated 000 masses switches Fig. 3. Mechanical testing of micro/minicomposites: (a) tensile specimens; (b) tensile device for minicomposite tests; and (c) apparatus for lifetime urements in air at high temperatures of SiC/SiC microcomposites, according to Ref. [ 9, 12, 26] interphases with elementary layer thicknesses of a few nm Vr=0.50, corresponding to that commonly observed for the can only be deposited (infiltrated) in a controlled manner by corresponding real nD-composites). The chemical analysis P-CVD/P-CVI [22, 23]. In a last step, the Sic (or carbon) of the FM-interfacial zone was performed according to two was deposited (infiltrated) by conventional CVD complementary techniques: (i) by Auger electron spectro- (CVI)in the apparatus used for the fabrication of real nd scopy (AES)depth profiles, recorded from fracture composites and under similar T-P conditions(typically at surfaces both on debonded fibers [see Fig. I(b) and on 00'C-1000C and under few kPa or few 10 kPa, depend- their trough in the matrix, and (i by parallel electron ing on the nature of the precursor). Hence, micro/mini energy loss spectroscopy(P-EELS)in a transmission elec tron microscope(TEM). Longitudinal thin foils of micro- ques closely related to those used for the corresponding real composites were prepared according to a technique which composites, with however two important differences: (i)the has been depicted elsewhere [11] processing time is much shorter and; (ii)the handling of the Micro-and mini-composites were tensile tested at room specimens requires specific care. Examples of micro/mini- temperature with tensile devices built in-house and composites are shown in Figs. I and 2 described elsewhere [25-27]. Microcomposites were tested The characterization of micro/mini composites requires specific experimental procedures owing to the small radial size of the specimens (a Nicalon/SiC microcomposite has an I Auger PHI 5590, from Physical Electronics PEELS 666-3K, from Gatan (USA). overall diameter of about 20 um for a fiber volume fraction, CM 30 ST, from Philips (NL)
interphases with elementary layer thicknesses of a few nm can only be deposited (infiltrated) in a controlled manner by P-CVD/P-CVI [22,23]. In a last step, the SiC (or carbon) matrix was deposited (infiltrated) by conventional CVD (CVI) in the apparatus used for the fabrication of real nD composites and under similar T–P conditions (typically at 9008C–10008C and under few kPa or few 10 kPa, depending on the nature of the precursor). Hence, micro/mini composites are fabricated according to processing techniques closely related to those used for the corresponding real composites, with however two important differences: (i) the processing time is much shorter and; (ii) the handling of the specimens requires specific care. Examples of micro/minicomposites are shown in Figs. 1 and 2. The characterization of micro/mini composites requires specific experimental procedures owing to the small radial size of the specimens (a Nicalon/SiC microcomposite has an overall diameter of about 20 mm for a fiber volume fraction, Vf 0.50, corresponding to that commonly observed for the corresponding real nD-composites). The chemical analysis of the FM-interfacial zone was performed according to two complementary techniques: (i) by Auger electron spectroscopy (AES)1 depth profiles, recorded from fracture surfaces both on debonded fibers [see Fig. 1(b)] and on their trough in the matrix, and (ii) by parallel electron energy loss spectroscopy (P-EELS)2 in a transmission electron microscope (TEM)3 . Longitudinal thin foils of microcomposites were prepared according to a technique which has been depicted elsewhere [11]. Micro- and mini-composites were tensile tested at room temperature with tensile devices built in-house and described elsewhere [25–27]. Microcomposites were tested R. Naslain et al. / Composites: Part A 30 (1999) 537–547 539 Fig. 3. Mechanical testing of micro/minicomposites: (a) tensile specimens; (b) tensile device for minicomposite tests; and (c) apparatus for lifetime measurements in air at high temperatures of SiC/SiC microcomposites, according to Ref. [9, 12, 26]. 1 Auger PHI 5590, from Physical Electronics. 2 PEELS 666-3 K, from Gatan (USA). 3 CM 30 ST, from Philips (NL)
R. Naslain et al/Composites: Part A 30(1999)537-547 MATRI 「atB①P Sputter time (min M 500nm nd TEM analysis of complex interphases in SiC/SiC microcomposites:(a)C (B) graded com ase; (b)AES depth profile (B)graded composition type A interphase(a being similar to A but wi sublayer V),(c)bright field TEM of a Nicalon/SiC with a type A'C(B)interphase, and(d) bright field TEM image of a Hi-Nicalon/SiC microcomposite with a(Py C-SiC)o multilayered interphase, according to Refs. [11, 12, 15, 23] according to a procedure similar to that one used for single elsewhere [5,6]. The FM-interfacial parameters were also filaments. A length of microcomposite was pasted with an assessed from push-in or push-through experiments epoxy cement on a paper holder(gauge length, L, ranging performed on polished cross-sections [14. Finally, the fail from 10 to 50 mm), as shown in Fig 3(a). The whole assem- ure surfaces were observed with a high-resolution scanning bly was then attached to the tensile tester grips and finally, microscope(HR-sEM) to identify the failure mode and to the paper holder was cut immediately before applying the calculate the in-situ failure stress of the fibers from mirror load. In a similar manner, minicomposites were attached radius measurements with the epoxy cement to two steel tubes(gauge lengt he effect of the environment, e.g. the ambient air, on the ranging from 50 to 75 mm)and strained at a constant mechanical behaviour and lifetime, was studied through speed(0.085% per min), the displacement being measured fatigue tests(static or cyclic)performed at high tempera either optically or with inductive transducers [Fig. 3(b)]. tures on either micro-or mini-composites The minicompo- Unloading-reloading hysteresis loops were systematically site specimens were prepared as described above for tests recorded in order to measure the Youngs modulus, E, and performed at room temperature, the epoxy cement being the residual permanent strain, ep, (at o=0), of the material replaced by a high temperature alumina-based cement. In as it is progressively damaged and to derive the FM-inter- the static(or cyclic)fatigue tests run on microcomposites, a facial parameters, i.e. the debond length, Id, the interfacial length of microcomposite was attached with the alumina- shear stress, Ti, and the debond energy, Ti, according to micromechanics-based models, which have been reported S4500, from Hitachi (Japan)
according to a procedure similar to that one used for single filaments. A length of microcomposite was pasted with an epoxy cement on a paper holder (gauge length, L, ranging from 10 to 50 mm), as shown in Fig. 3(a). The whole assembly was then attached to the tensile tester grips and finally, the paper holder was cut immediately before applying the load. In a similar manner, minicomposites were attached with the epoxy cement to two steel tubes (gauge length ranging from 50 to 75 mm) and strained at a constant speed (0.085% per min), the displacement being measured either optically or with inductive transducers [Fig. 3(b)]. Unloading–reloading hysteresis loops were systematically recorded in order to measure the Young’s modulus, E, and the residual permanent strain, ep, (at s 0), of the material as it is progressively damaged and to derive the FM-interfacial parameters, i.e. the debond length, ld, the interfacial shear stress, ti, and the debond energy, Gi, according to micromechanics-based models, which have been reported elsewhere [5,6]. The FM-interfacial parameters were also assessed from push-in or push-through experiments performed on polished cross-sections [14]. Finally, the failure surfaces were observed with a high-resolution scanning microscope (HR-SEM)4 to identify the failure mode and to calculate the in-situ failure stress of the fibers from mirror radius measurements. The effect of the environment, e.g. the ambient air, on the mechanical behaviour and lifetime, was studied through fatigue tests (static or cyclic) performed at high temperatures on either micro- or mini-composites. The minicomposite specimens were prepared as described above for tests performed at room temperature, the epoxy cement being replaced by a high temperature alumina-based cement. In the static (or cyclic) fatigue tests run on microcomposites, a length of microcomposite was attached with the alumina- 540 R. Naslain et al. / Composites: Part A 30 (1999) 537–547 Fig. 4. AES and TEM analysis of complex interphases in SiC/SiC microcomposites: (a) C (B) graded composition type A 0 interphase; (b) AES depth profile analysis of a C (B) graded composition type A interphase (A being similar to A 0 but without sublayer V); (c) bright field TEM image of a Nicalon/SiC microcomposite with a type A 0 C (B) interphase; and (d) bright field TEM image of a Hi-Nicalon/SiC microcomposite with a (PyC-SiC)10 multilayered interphase, according to Refs. [11,12,15,23]. 4 S 4500, from Hitachi (Japan)
R. Naslain et al/Composites: Part A 30(1999)537-547 based cement to two 100 um SIC CVD filaments, acting elementary thicknesses are of the order of 100 nm as tensile rods, and the lifetime was measured with a timer as expected; and (ii) a matrix mode-I crack, which has been shown in Fig 3(c) deflected in mode Il(parallel to the fiber surface) near th boundary between sublayers Il and ll [11, 12]. The analys 3. Results and discussion of complex multilayered interphases becomes more difficult when the thickness of the elementary sublayer is extremely The objective of the present contribution being to show small, typically a few nm. Under such conditions, the AES microprobe resolution may be insuficient and the TEM/P- the potential and limits of the micro/mini composites EEls analysis the only appropriate technique. As an exam- approach, rather than to give a detailed analysis of the mate- als, a few examples will be presented and discussed to ple, Fig. 4(d)shows a TEM-image recorded from a lon illustrate how this approach can be used to assess the struc. itudinal thin foil of a Hi-Nicalon/(Py C-SiC)Sic tural and mechanical behaviour of selected composites as microcomposite failed under tensile loading. First, the 10 well as the effect of an oxidizing atmosphere PyC-SiC sequences deposited by P-CVD(each comprising a 20 nm thick pyrocarbon and a 30 nm thick SiC +Cnano 3.1. Chemical and structural analysis crystalline sublayers)are clearly apparent. Second, the TEM-image also shows a mode-I matrix microcrack The most complex region in a CMC is the FM-interfacial which has been multideflected parallel to the fiber surface zone CMCs display a non-brittle mechanical behaviour in the interphase [23]. When necessary, selected area when the FM-bonding is not too strong. The control of the diffraction(SAD), lattice fringe(LF) image and P-EELS FM-bonding is achieved during processing, via the use of a elementary analysis are used to get an insight in the struc- thin layer of a compliant material with a low shear strength ture and composition of the FM-interfacial zone in micro/ red to as the interphase. It has been suggested that minI composites appropriate interphase materials might be those with a layered crystal structure(pyrocarbon or hex-BN)or a 3. 2. Mechanical behaviour at room temperature ayered microstructure, such as the(PyC-SiC )n interphase [22] with typically n= 1-10. The main function of the As shown in Fig. 5, the tensile stress-strain curves fo interphase is to act as mechanical fuse, i.e. to deflect the micro- and minl-composites display the same general matrix microcracks parallel to the fiber surface. The chemi features as those reported for their multidirectional counter cal composition, morphology and structure of highly engi- parts. Beyond the proportional limit, both model composites peered interphases, e.g. multilayered interphases, can be undergo damage phenomena, I.e. mainly matrix multiple extremely complex. However, their analysis can be cracking and fiber debonding, which are responsible for performed in a rather straightforward manner on either the non-linear feature of the stress-strain behaviour. As micro-or mini-composite specimens, provided appropriate strain increases, the model composites are more damaged experimental procedures are used, as already said in the with the result that: (1) the stiffness of the composites, preceding section. An example of such an analysis is assessed through secant modulus measurements, decreases shown in Fig. 4, for two SiC/SiC microcomposites with (1i) simultaneously, the width, 84, and the area,S, of the different fibers and interphases. In the Nicalon/C( B)/Sic unloading-reloading hysteresis loops increase and finally microcomposite, the interphase consists of 4(or 5)boron- (ii) the permanent residual strain(at O =O),Ep, increases doped 100 nm pyrocarbon sublayers. The first sublayer All these features depend on the intrinsic failure properties deposited on the fiber surface is a film of pure anisotropic of the brittle matrix and the characteristics of the FM-bond- pyrocarbon(deposited from propane) whereas the followi ing and are thus observed for both model and real compo- sublayers exhibit a boron concentration which increases sites. However, there are also differences which are related when moving towards the matrix(the gaseous precur to the nature of the fiber architecture and processing consid- being C3 Hg-BCI3-H2). Boron increases the anisotropy of erations. As an example, in the 2D-Nicalon/PyC/Sic pyrocarbon at low doping levels(with a maximum effect composites processed with a relatively strong FM-bonding for =8 at. %B; sublayer II)and improves its oxidation and exhibiting a high failure strain, the presence of three resistance at high doping levels(sublayers Ill-V). The families of matrix cracks has been reported, which are omposition gradient in the multilayered interphase is successively formed as the strain increases. The first family clearly apparent from the AES depth profiles, which have consists of cracks initiated at the large residual intertow een recorded from the coated fiber surface prior to the sic pores left by the CVI-process. The second family comprises matrix deposition [Fig 4(b)]. Furthermore, the TEM-image cracks formed in the transverse tows( those oriented at 90 recorded on a longitudinal thin foil of a microcomposite to the load assumed to be applied along the O longitudinal failed under tensile loading [ Fig 4(c) 1, shows:(i) the inter- tows). Finally, the third is made of transverse cracks present nal structure of the interphase i.e. five sublayers whose in the O tows. It has been shown that the longitudinal tows behave as physical entities, and dictate the mechanical beha SIGMA, Germany viour of the composite once the first and the second familie
based cement to two 100 mm SiC CVD filaments5 , acting as tensile rods, and the lifetime was measured with a timer as shown in Fig. 3(c). 3. Results and discussion The objective of the present contribution being to show the potential and limits of the micro/mini composites approach, rather than to give a detailed analysis of the materials, a few examples will be presented and discussed to illustrate how this approach can be used to assess the structural and mechanical behaviour of selected composites as well as the effect of an oxidizing atmosphere. 3.1. Chemical and structural analysis The most complex region in a CMC is the FM-interfacial zone. CMCs display a non-brittle mechanical behaviour when the FM-bonding is not too strong. The control of the FM-bonding is achieved during processing, via the use of a thin layer of a compliant material with a low shear strength referred to as the interphase. It has been suggested that appropriate interphase materials might be those with a layered crystal structure (pyrocarbon or hex-BN) or a layered microstructure, such as the (PyC-SiC)n interphase [22] with typically n 1–10. The main function of the interphase is to act as mechanical fuse, i.e. to deflect the matrix microcracks parallel to the fiber surface. The chemical composition, morphology and structure of highly engineered interphases, e.g. multilayered interphases, can be extremely complex. However, their analysis can be performed in a rather straightforward manner on either micro- or mini-composite specimens, provided appropriate experimental procedures are used, as already said in the preceding section. An example of such an analysis is shown in Fig. 4, for two SiC/SiC microcomposites with different fibers and interphases. In the Nicalon/C (B) /SiC microcomposite, the interphase consists of 4 (or 5) borondoped 100 nm pyrocarbon sublayers. The first sublayer deposited on the fiber surface is a film of pure anisotropic pyrocarbon (deposited from propane) whereas the following sublayers exhibit a boron concentration which increases when moving towards the matrix (the gaseous precursor being C3H8-BCl3-H2). Boron increases the anisotropy of pyrocarbon at low doping levels (with a maximum effect for < 8 at.% B; sublayer II) and improves its oxidation resistance at high doping levels (sublayers III–V). The composition gradient in the multilayered interphase is clearly apparent from the AES depth profiles, which have been recorded from the coated fiber surface prior to the SiCmatrix deposition [Fig. 4(b)]. Furthermore, the TEM-image recorded on a longitudinal thin foil of a microcomposite failed under tensile loading [Fig. 4(c)], shows: (i) the internal structure of the interphase i.e. five sublayers whose elementary thicknesses are of the order of 100 nm as expected; and (ii) a matrix mode-I crack, which has been deflected in mode II (parallel to the fiber surface) near the boundary between sublayers II and III [11,12]. The analysis of complex multilayered interphases becomes more difficult when the thickness of the elementary sublayer is extremely small, typically a few nm. Under such conditions, the AES microprobe resolution may be insufficient and the TEM/PEELS analysis the only appropriate technique. As an example, Fig. 4(d) shows a TEM-image recorded from a longitudinal thin foil of a Hi-Nicalon/(PyC-SiC)10/SiC microcomposite failed under tensile loading. First, the 10 PyC-SiC sequences deposited by P-CVD (each comprising a 20 nm thick pyrocarbon and a 30 nm thick SiC 1 C nanocrystalline sublayers) are clearly apparent. Second, the TEM-image also shows a mode-I matrix microcrack which has been multideflected parallel to the fiber surface in the interphase [23]. When necessary, selected area diffraction (SAD), lattice fringe (LF) image and P-EELS elementary analysis are used to get an insight in the structure and composition of the FM-interfacial zone in micro/ mini composites. 3.2. Mechanical behaviour at room temperature As shown in Fig. 5, the tensile stress–strain curves for micro- and mini-composites display the same general features as those reported for their multidirectional counterparts. Beyond the proportional limit, both model composites undergo damage phenomena, i.e. mainly matrix multiple cracking and fiber debonding, which are responsible for the non-linear feature of the stress–strain behaviour. As strain increases, the model composites are more damaged with the result that: (i) the stiffness of the composites, assessed through secant modulus measurements, decreases; (ii) simultaneously, the width, dD, and the area, S, of the unloading–reloading hysteresis loops increase and finally; (iii) the permanent residual strain (at s 0), ep , increases. All these features depend on the intrinsic failure properties of the brittle matrix and the characteristics of the FM-bonding and are thus observed for both model and real composites. However, there are also differences which are related to the nature of the fiber architecture and processing considerations. As an example, in the 2D-Nicalon/PyC/SiC composites processed with a relatively strong FM-bonding and exhibiting a high failure strain, the presence of three families of matrix cracks has been reported, which are successively formed as the strain increases. The first family consists of cracks initiated at the large residual intertow pores left by the CVI-process. The second family comprises cracks formed in the transverse tows (those oriented at 908 to the load assumed to be applied along the 08 longitudinal tows). Finally, the third is made of transverse cracks present in the 08 tows. It has been shown that the longitudinal tows behave as physical entities, and dictate the mechanical behaviour of the composite once the first and the second families R. Naslain et al. / Composites: Part A 30 (1999) 537–547 541 5 SIGMA, Germany