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 familiebased 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 mate￾rials, a few examples will be presented and discussed to illustrate how this approach can be used to assess the struc￾tural 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 ˆ 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 chemi￾cal composition, morphology and structure of highly engi￾neered 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) boron￾doped 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 SiC￾matrix 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 inter￾nal 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/P￾EELS analysis the only appropriate technique. As an exam￾ple, Fig. 4(d) shows a TEM-image recorded from a long￾itudinal 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 nano￾crystalline 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 struc￾ture 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 counter￾parts. 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 ˆ 0), ep , increases. All these features depend on the intrinsic failure properties of the brittle matrix and the characteristics of the FM-bond￾ing and are thus observed for both model and real compo￾sites. However, there are also differences which are related to the nature of the fiber architecture and processing consid￾erations. 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 beha￾viour 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