E噩≈3S Journal of the European Ceramic Society 20(2000)1-13 TEM structure of(PyC/SiC)n multilayered interphases in SiC/SiC composites S. Bertrand * C. Droillard.R. Pailler. X. Bourrat.R. Naslain Laboratoire des Composites Thermostructuraux, UMR 580/ CNRS-SEP/SNECMA-UBl, Universite bordeaux, 1-3 allee de la boetie, Received 28 January 1999; accepted 14 March 1999 Two generations of multilayered interphases, composed of carbon and silicon carbide, have been developed to act as a mechan- ical fuse in SiC/Sic composites with improved oxidation resistance. Pyrocarbon is an ideal interfacial material, from the mechanical point of view, whereas Sic has a good oxidation resistance. In the multilayered interphase, the carbon mechanical fuse is split into hin sublayers, each being protected against oxidation by the neighbouring Sic-based glass former layers. A first generation of multilayers as synthesised by means of isobaric-CVI with sublayers with micrometric thickness. Then, in order to push forward the concept, pressure pulsed-CVI was involved to deposit nanometric scale sublayers. In this work, transmission electron microscopy was developed to characterise the two generations of materials. The microstructure of the layers and the influence of the fibrous preforms on the structure of the layers were studied. Examinations were then performed on the loaded samples and damaging mode haracterised at nanometric scale. C 1999 Elsevier Science Ltd. All rights reserved Keywords: Composites; Electron microscopy: Interphase: SiC; Carbon 1. ntroduction types of interphase involving alternating thin layers of two different materials have been suggested: the laminar It is now well established that the mechanical beha ceramics, 6-10 viour of ceramic matrix composites(CMCs) with con- A breakthrough was achieved by Droillard et al I1, I tinuous fibre reinforcement depends not only on the demonstrating that(PyC/SiC)n multilayered material intrinsic properties of the fibre and the matrix, but also in 2D woven Nicalon SiC composites, behaves as an n the fibre-matrix bonding. -To control the strength efficient interfacial materials, but only if their bonding of the fibre-matrix bonding in CMCs, an additional to the fibre surface is reinforced. Fig. 1, published in a phase referred to as the interphase is used which serves previous paper I shows the tensile tests realised on 2D as a compliant layer between the fibre and matrix. Two Nicalon/SiC composites with different multilayered main functions are devoted to the interphase: first, load combinations. All the materials could be grouped into transfer between matrix and reinforcement and sec- two distinct families: (i) materials reinforced with ondly, control of the crack deflection at the interface. 4 untreated fibres have a weak fibre bonding and are The interphase is deposited on the fibre surface prior to characterised by a relatively low strength and a low interphase materials are pyrocarbon(PyC) and boron possess a stronger interface and are characterised ba the deposition of the matrix. The most commonly used toughness, whereas (i) materials with treated fibre nitride(BN). However, both of them are not stable high strength and a high toughness. As a result, when under oxidising conditions at high temperatures. New stronger interfaces were introduced, strength and concepts have been proposed to produce interphase that toughness were increased; in the mean time more than have both oxidation resistance and mechanical proper- 50% of the carbon was removed from the interfacial ties required to yield tough composites. 4.5 Also, new zone. In contrast, when the interface was weak, only the first carbon sublayer was involved in the fracture mechanism and ultimate performances remained s Corresponding author 0955-2219/99/.see front C 1999 Elsevier Science Ltd. All rights reserved PII:S0955-2219(99)00
TEM structure of (PyC/SiC)n multilayered interphases in SiC/SiC composites S. Bertrand*, C. Droillard, R. Pailler, X. Bourrat, R. Naslain Laboratoire des Composites Thermostructuraux, UMR 5801 CNRS-SEP/SNECMA-UB1, Universite Bordeaux, 1±3 alleÂe de la BoeÂtie, 33 600 Pessac, France Received 28 January 1999; accepted 14 March 1999 Abstract Two generations of multilayered interphases, composed of carbon and silicon carbide, have been developed to act as a mechanical fuse in SiC/SiC composites with improved oxidation resistance. Pyrocarbon is an ideal interfacial material, from the mechanical point of view, whereas SiC has a good oxidation resistance. In the multilayered interphase, the carbon mechanical fuse is split into thin sublayers, each being protected against oxidation by the neighbouring SiC-based glass former layers. A ®rst generation of multilayers as synthesised by means of isobaric-CVI with sublayers with micrometric thickness. Then, in order to push forward the concept, pressure pulsed-CVI was involved to deposit nanometric scale sublayers. In this work, transmission electron microscopy was developed to characterise the two generations of materials. The microstructure of the layers and the in¯uence of the ®brous preforms on the structure of the layers were studied. Examinations were then performed on the loaded samples and damaging mode characterised at nanometric scale. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: Composites; Electron microscopy; Interphase; SiC; Carbon 1. Introduction It is now well established that the mechanical behaviour of ceramic matrix composites (CMCs) with continuous ®bre reinforcement depends not only on the intrinsic properties of the ®bre and the matrix, but also on the ®bre±matrix bonding.1±3 To control the strength of the ®bre±matrix bonding in CMCs, an additional phase referred to as the interphase is used which serves as a compliant layer between the ®bre and matrix. Two main functions are devoted to the interphase: ®rst, load transfer between matrix and reinforcement and secondly, control of the crack de¯ection at the interface.4 The interphase is deposited on the ®bre surface prior to the deposition of the matrix. The most commonly used interphase materials are pyrocarbon (PyC) and boron nitride (BN). However, both of them are not stable under oxidising conditions at high temperatures. New concepts have been proposed to produce interphase that have both oxidation resistance and mechanical properties required to yield tough composites.4,5 Also, new types of interphase involving alternating thin layers of two dierent materials have been suggested: the laminar ceramics.6±10 A breakthrough was achieved by Droillard et al.11,12 demonstrating that (PyC/SiC)n multilayered materials, in 2D woven Nicalon/SiC composites, behaves as an ecient interfacial materials, but only if their bonding to the ®bre surface is reinforced. Fig. 1, published in a previous paper,11 shows the tensile tests realised on 2D Nicalon/SiC composites with dierent multilayered combinations. All the materials could be grouped into two distinct families: (i) materials reinforced with untreated ®bres have a weak ®bre bonding and are characterised by a relatively low strength and a low toughness, whereas (ii) materials with treated ®bres possess a stronger interface and are characterised by a high strength and a high toughness. As a result, when stronger interfaces were introduced, strength and toughness were increased; in the mean time more than 50% of the carbon was removed from the interfacial zone. In contrast, when the interface was weak, only the ®rst carbon sublayer was involved in the fracture mechanism11 and ultimate performances remained low. 0955-2219/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0955-2219(99)00086-2 Journal of the European Ceramic Society 20 (2000) 1±13 * Corresponding author
400 B H 300 OTHON INPRPACP a25星星4 K C 200 0 WIX TERFACE 02 0.4 0.6 Longitudinal tensile strain(c) Fig. I. Tensile stress-strain curves obtained for 2D-SiCSiC composites with various multilayered interphases: B, D, H and L fabricated with treated Nicalon fibre; A, C, G and K fabricated with as-received Nicalon fibre(according to Refs [4] and [liD Pasquier3 has investigated the potentialities of multi- tooling. The components of the multilayered(Py C/SiC) layered(PyC/SiC)n interphases at the micrometer scale in interphases and the Sic-matrix were infiltrated within the Nicalon/SiC composites in terms of oxidation resistance porous fibre preforms, according to the isothermal/iso- Then, Heurtevent4 has developed the nanoscale multi- baric chemical vapour infiltration(I-CVI)process, which layered(PyC/SiC)n interphases in Hi-Nicalon/ SiC micro- has been described elsewhere. 6-18 Pyrocarbon and silicon composites and studied their behaviour at high carbide were deposited from propane C3Hs and methyl temperatures in oxidative conditions trichlorosilane(MTS) CH3siCl3/H2, respectively, accord The aim of the present paper is to char ng to the following overall equations: tructure of micro- and nano-scaled(Py C/SiC)n multi layered interphases. The particular aspect related to the CH3 SiC3(e) SiC+3HClgy treatment is fully described in a companion paper, s ace first interface and the influence of the fibre surf C3Hxg)→3C(+4H2g 2. Experimental procedure in a hot-wall chamber (internal diameter: 130 mm; 2.1. 2D-SiC/SiC materials obtained by 1-CVI cooled r f. coil (maximum temperature capability c1500.C). The apparatus has been designed to work The 2D-SiC/SiC composites were prepared as rectan- under reduced pressures(0.5<P<10 kPa), the total gular plates (130x 100x5 mm )from 2D-preforms con- pressure being maintained at a constant value with a sisting of stacks of Nicalon fabrics(NLM 202 ceramic pressure sensor(type 127A from MKS)and a pressure grade from Nippon Carbon Company Ltd, Tokyo, regulator(type 252A from MKS). Mass flowmeters were Japan)maintained pressed together with a graphite used to measure the flowrates of the various gaseous
Pasquier13 has investigated the potentialities of multilayered (PyC/SiC)n interphases at the micrometer scale in Nicalon/SiC composites in terms of oxidation resistance. Then, Heurtevent14 has developed the nanoscale multilayered (PyC/SiC)n interphases in Hi-Nicalon/SiC microcomposites and studied their behaviour at high temperatures in oxidative conditions. The aim of the present paper is to characterise the structure of micro- and nano-scaled (PyC/SiC)n multilayered interphases. The particular aspect related to the ®rst interface and the in¯uence of the ®bre surface treatment is fully described in a companion paper.15 2. Experimental procedure 2.1. 2D-SiC/SiC materials obtained by I-CVI The 2D-SiC/SiC composites were prepared as rectangular plates (1301005 mm3 ) from 2D-preforms consisting of stacks of Nicalon fabrics (NLM 202 ceramic grade from Nippon Carbon Company Ltd., Tokyo, Japan) maintained pressed together with a graphite tooling. The components of the multilayered (PyC/SiC)n interphases and the SiC-matrix were in®ltrated within the porous ®bre preforms, according to the isothermal/isobaric chemical vapour in®ltration (I-CVI) process, which has been described elsewhere.16±18 Pyrocarbon and silicon carbide were deposited from propane C3H8 and methyltrichlorosilane (MTS) CH3SiCl3/H2, respectively, according to the following overall equations: CH3SiC3 g ! H2 SiC s 3HCl g 1 C3H8 g ! 3C s 4H2 g 2 in a hot-wall chamber (internal diameter: 130 mm; height: 250 mm) heated isothermally with a watercooled r.f. coil (maximum temperature capability: 1500C). The apparatus has been designed to work under reduced pressures (0.5<P<10 kPa), the total pressure being maintained at a constant value with a pressure sensor (type 127A from MKS) and a pressure regulator (type 252A from MKS). Mass ¯owmeters were used to measure the ¯owrates of the various gaseous Fig. 1. Tensile stress±strain curves obtained for 2D-SiC/SiC composites with various multilayered interphases: B, D, H and L fabricated with treated Nicalon ®bre; A, C, G and K fabricated with as-received Nicalon ®bre (according to Refs.[4] and [11]). 2 S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13
species: propane (Qc Hs), hydrogen(QH and MTs sublayers has been limited to 0. 1 um or less for oxidation (QMS) resistance considerations(which will be not developed MTS, being a liquid at room temperature, is evapo- here). rated in a boiler set in a drying oven and then, the MTS(g is mixed with hydrogen before being injected in 2. 2. Pressure pulsed-CVl. nanometric scale multilayers the infiltration chamber. The composition of the gas phase used in the infiltration of Sic is characterised by a The nanometric scale multilayered interphases have dilution a-ratio defined as been fabricated by pressure pulsed CVD/CVI (P-CVD/ P-CVI) process. The apparatus used for the fabrication PH, QH of the materials and the experimental conditions have been described in another article 20 In this process, the operating pressure is pi acting gas is allowed by an upper inlet pneumatic valve, up to where Pi and Qi (with 1= H2 or MTS)are the partial the operating pressure. Then the furnace is closed dur- pressure and gas flowrate of species i, respectively. The ing a residence time, Ir and, finally, it is evacuated experiments have been carried out under conditions through an outlet pneumatic valve, and cooled traps by typical of the I-CVi process, which have been discussed using a rotary pump. A computer is used to monitor elsewhere valves'opening and closing, safety devices and the total Two different series of 2D-SiC/SiC composites have amount of pulses been prepared. In the first series, the Nicalon fabrics Hi-Nicalon bundles(from Nippon Carbon Company were used as-received (materials A, C, G, and K in Ltd, Tokyo, Japan), were used for the fabrication of Table 1)whereas, in the second series(materials B, D, SiC/Sic minicomposites(a minicomposite is a model H, and L), the Nicalon fabrics have received a treatment ID composite with one single fibre tow). In order to (proprietary treatment performed by SEP, Bordeaux) change the interfacial bonding strength two series of prior to the infiltration of the multilayered interphase, reinforcement were systematically utilised: in the first in order to improve the fibre-matrix bonding. series, tows were used as-received (i.e. non treated The nature of the various multilayered interphases fibres) whereas, in the second series, the tows were pre- deposited by I-CVI on the fibre surface is shown in viously submitted to a treatment (so-called treated Table 1. The interphases exhibit the following features: fibres) performed at the Laboratory, prior to the infil (i the first sublayer (i.e. that in contact with the fibre tration (or deposition) of the multilayered interphase surface)is always a pyrocarbon sublayer and the inter- The nature of the various multilayered interphase is facial sequence can be written as(PyC/SiC)n, (ii)when hown in table 2 n>1, the Sic of the last sequence is that of the matrix, (ii) the overall thickness of the multilayered interphase 2.3. Microstructural characterisation is constant and equal to 0.5 um and (iv) the thicknesses of the C and Sic sublayers are either maintained con- Microstructure of the multil stant(materials evolutive(materials K, L) ssessed by Transmission Electron Microscopy(TEM within the interphase. Finally, the thickness of the Pyc CM30ST/PEELS from Philips). TEM analyses were performed on sample cross-sections, perpendicular to the axis of the fibres as well as longitudinal sections Table TEM specimen sampling has been previously detailed in Material references and nature of the multilayered interphases of the 2D Nicalon/ SiC composites, processed by I-CVI a companion paper Optical microscopy in polarised light(MeF3 from Materials Nature of Nature of the C-SiC sequence in Reichert-Jung) was used to measure the pyrocarbon the fabrics the interphase and thickness(in um) anisotropy following the extinction angle technique F/C/SiC/C/Md (Ae), fully described elsewhere. 21 Ae was observed to 0.10.30.1 fall between I2°and14° corresponding to sm laminar(SL) and rough laminar(RL) pyrocarbon 0.10.10.10.10.1 X-Ray Diffraction (XRD, Siemens D5000)analysis 0.050.10050.10.050.10.05 was performed in order to evaluate the apparent crys tallite size in the [111] crystallographic direction (Lul) 0.050050.050.10.050.150.05 of the Sic grains by means of the Scherrer equation NT: not treated b T treated Auger Electron Spectroscopy (AES) microprobe equipped with an Ar+ sputtering gun(VG 310F)was used to record composition-depth profiles
species: propane (QC3H8 ), hydrogen (QH2 ) and MTS (QMTS). MTS, being a liquid at room temperature, is evaporated in a boiler set in a drying oven and then, the MTS g is mixed with hydrogen before being injected in the in®ltration chamber. The composition of the gas phase used in the in®ltration of SiC is characterised by a dilution -ratio de®ned as: PH2 PMTS QH2 QMTS 3 where Pi and Qi (with i H2 or MTS) are the partial pressure and gas ¯owrate of species i, respectively. The experiments have been carried out under conditions typical of the I-CVI process, which have been discussed elsewhere.16±18 Two dierent series of 2D-SiC/SiC composites have been prepared. In the ®rst series, the Nicalon fabrics were used as-received (materials A, C, G, and K in Table 1) whereas, in the second series (materials B, D, H, and L), the Nicalon fabrics have received a treatment (proprietary treatment performed by SEP, Bordeaux) prior to the in®ltration of the multilayered interphase, in order to improve the ®bre±matrix bonding.19 The nature of the various multilayered interphases deposited by I-CVI on the ®bre surface is shown in Table 1. The interphases exhibit the following features: (i) the ®rst sublayer (i.e. that in contact with the ®bre surface) is always a pyrocarbon sublayer and the interfacial sequence can be written as (PyC/SiC)n, (ii) when n > 1, the SiC of the last sequence is that of the matrix, (iii) the overall thickness of the multilayered interphase is constant and equal to 0.5 mm and (iv) the thicknesses of the C and SiC sublayers are either maintained constant (materials G, H) or evolutive (materials K, L) within the interphase. Finally, the thickness of the PyC sublayers has been limited to 0.1 mm or less for oxidation resistance considerations (which will be not developed here). 2.2. Pressure pulsed-CVI: nanometric scale multilayers The nanometric scale multilayered interphases have been fabricated by pressure pulsed CVD/CVI (P-CVD/ P-CVI) process. The apparatus used for the fabrication of the materials and the experimental conditions have been described in another article.20 In this process, the operating pressure is pulsed. First, admission of reacting gas is allowed by an upper inlet pneumatic valve, up to the operating pressure. Then the furnace is closed during a residence time, tr and, ®nally, it is evacuated through an outlet pneumatic valve, and cooled traps by using a rotary pump. A computer is used to monitor valves' opening and closing, safety devices and the total amount of pulses. Hi-Nicalon bundles (from Nippon Carbon Company Ltd., Tokyo, Japan), were used for the fabrication of SiC/SiC minicomposites (a minicomposite is a model 1D composite with one single ®bre tow). In order to change the interfacial bonding strength two series of reinforcement were systematically utilised: in the ®rst series, tows were used as-received (i.e. non treated ®bres) whereas, in the second series, the tows were previously submitted to a treatment (so-called treated ®bres) performed at the Laboratory, prior to the in®ltration (or deposition) of the multilayered interphase. The nature of the various multilayered interphase is shown in Table 2. 2.3. Microstructural characterisation Microstructure of the multilayers was essentially assessed by Transmission Electron Microscopy (TEM, CM30ST/PEELS from Philips). TEM analyses were performed on sample cross-sections, perpendicular to the axis of the ®bres as well as longitudinal sections. TEM specimen sampling has been previously detailed in a companion paper.15 Optical microscopy in polarised light (MeF3 from Reichert-Jung) was used to measure the pyrocarbon anisotropy following the extinction angle technique (Ae), fully described elsewhere.21 Ae was observed to fall between 12 and 14 corresponding to smooth laminar (SL) and rough laminar (RL) pyrocarbon. X-Ray Diraction (XRD, Siemens D5000) analysis was performed in order to evaluate the apparent crystallite size in the [111] crystallographic direction (L111) of the SiC grains by means of the Scherrer equation (k=0.9). Auger Electron Spectroscopy (AES) microprobe equipped with an Ar+ sputtering gun (VG 310F) was used to record composition-depth pro®les. Table 1 Material references and nature of the multilayered interphases of the 2D Nicalon/SiC composites, processed by I-CVI Materials Nature of the fabrics Nature of the C-SiC sequence in the interphase and thickness (in mm) A NTa Fc /C/SiC/C/Md B Tb 0.1 0.3 0.1 C NT F/C/C/SiC/M D T 0.1 0.1 0.1 0.1 0.1 G NT F/C/SiC/C/SiC/C/M H T 0.05 0.1 0.05 0.1 0.05 0.1 0.05 K NT F/C/SiC/C/SiC/C/M L T 0.05 0.05 0.05 0.1 0.05 0.15 0.05 a NT: not treated. b T: treated. c F: ®bre. d M: matrix. S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13 3
Fracture surfaces were examined by Scanning Electron 3.1. Pyrocarbon nanostructure Microscopy(FEG-SEM Hitachi $4500)at low voltage (3kv) Generally speaking, pyrocarbon resulting from the cracking of propane under the I-CVI conditions was characterised by a large value of the Lz-parameter, 3. Results--micrometer-scale multilayers as processed strong anisotropy and a low porosit by I-Cvi the lateral size of the aromatic carbon sheet in the tur bostratic stack as measured by Hr-tEm). Similar fea When seen in cross section, multilayers deposited by tures have been also reported for pyrocarbons resulting means of I-CVI exhibited rough and discontinuous from the cracking of propylene C3.23 sublayers(Fig. 2). The inset shows a low magnification The first pyrocarbon sublayer growth occurred of interphase"G" constituted by seven sublayers infil- directly onto the fibre surface whose composition was trated in an as-received 2d Nicalon nlm 202-based different for the two series of materials considered here preform. This sequenced ceramic material appears The bonding of carbon onto the fibre was seen to con- gh and disrupted. a close inspection at higher mag- trol the nature of the composite interface. 5Then, all nification revealed that carbon sublayers were system the subsequent pyrocarbon sublayers grew onto surfaces atically continuous, and that disruptions, when present, made of pure, well crystallised Sic which exhibited were related to the Sic sublayer crystallinity some roughness at the nanometric scale. As shown in Fig. 3, the Pyc deposit first filled the concave parts (formed by adjacent cone-like SiC crystals) of the SiC substrate. Then, at a distance, the PyC aromatic layers Material references and nature of the multilayered interphases of the tended to deposit parallel to the mean surface of the Hi-Nicalon/SiC minicomposites, processed by P-CVI coated fibre and exhibited a pronounced anisotropy Materials Nature of Nature of the C-SiC sequence in The analysis of the first carbon layers has not shown(on ne tows the interphase and thickness(in nm) the basis of the TEM images) any significant difference in the carbon organisation depending on the nature of 2050 the sublying Sic crystals F/(PyC/SiC)o/M 330 NT: non treated b T: treated Pyc sic 2 c2 10 naterial G (TEM contrasted brightfield): undulation of the layers related to the crystal finity of Sic. Inset is a low magnification of an equivalent area(same Fig. 3. Growth of the first pyrocarbon layers onto a well crystallised SiC surface: smoothing effect of carbon (high resolution TEM)
Fracture surfaces were examined by Scanning Electron Microscopy (FEG-SEM Hitachi S4500) at low voltage (3 kV). 3. ResultsÐmicrometer-scale multilayers as processed by I-CVI When seen in cross section, multilayers deposited by means of I-CVI exhibited rough and discontinuous sublayers (Fig. 2). The inset shows a low magni®cation of interphase ``G'' constituted by seven sublayers in®ltrated in an as-received 2D Nicalon NLM 202-based preform. This sequenced ceramic material appears rough and disrupted. A close inspection at higher magni®cation revealed that carbon sublayers were systematically continuous, and that disruptions, when present, were related to the SiC sublayer crystallinity. 3.1. Pyrocarbon nanostructure Generally speaking, pyrocarbon resulting from the cracking of propane under the I-CVI conditions was characterised by a large value of the L2-parameter, a strong anisotropy and a low porosity22 (L2 characterises the lateral size of the aromatic carbon sheet in the turbostratic stack as measured by HR-TEM). Similar features have been also reported for pyrocarbons resulting from the cracking of propylene C3H6. 23 The ®rst pyrocarbon sublayer growth occurred directly onto the ®bre surface whose composition was dierent for the two series of materials considered here. The bonding of carbon onto the ®bre was seen to control the nature of the composite interface.15 Then, all the subsequent pyrocarbon sublayers grew onto surfaces made of pure, well crystallised SiC which exhibited some roughness at the nanometric scale. As shown in Fig. 3, the PyC deposit ®rst ®lled the concave parts (formed by adjacent cone-like SiC crystals) of the SiCsubstrate. Then, at a distance, the PyC aromatic layers tended to deposit parallel to the mean surface of the coated ®bre and exhibited a pronounced anisotropy. The analysis of the ®rst carbon layers has not shown (on the basis of the TEM images) any signi®cant dierence in the carbon organisation depending on the nature of the sublying SiC crystals. Table 2 Material references and nature of the multilayered interphases of the Hi-Nicalon/SiC minicomposites, processed by P-CVI Materials Nature of the tows Nature of the C-SiC sequence in the interphase and thickness (in nm) 15 NTa Fc /(PyC/SiC)10/Md 54 Tb 20 50 45 NT F/(PyC/SiC)10/M 3 30 a NT: non treated. b T: treated. c F: ®bre. d M: matrix. Fig. 2. Cross-section of the interfacial sequence in material G (TEM contrasted bright®eld): undulation of the layers related to the crystallinity of SiC. Inset is a low magni®cation of an equivalent area (same technique). Fig. 3. Growth of the ®rst pyrocarbon layers onto a well crystallised SiC surface: smoothing eect of carbon (high resolution TEM). 4 S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13
3.2. Nanostructure of the Sic-sublayers The first Sic sublayer was observed to be often dis- continuous, particularly when its thickness was low (i.e. The Sic present in the multilayered(PyC/SiC)n inter- 0.1 um), as shown in Fig. 2. As a result, "mechanical phase was always deposited onto pyrocarbon surfaces, bridges"were formed under such conditions, between the roughly flat, as supported by a comparison of the pyr- sublayers(inset in Fig. 2). Obviously, an optimising of the ocarbon C(002) lattice fringes TEM images (recorded at nucleation/growth processes should be further carried out about the same magnification) shown in Figs. 3 and 4. Fig. in order to achieve thinner and smoother sublayer 4 shows that there is probably e relation between the orientation of the carbon aromatic planes in the substrate 3.3. Multiple interfacing and that of the crystals in the Sic-deposit. Meanwhile, no precise relationship was found for the preferred growth The number of interfaces in the(PyC/SiC)n multi directions of SiC (i.e. 1 ll for the cubic B modification and layered interphase increases as n is raised: up to 8 when 00. 1 for the hexagonal a modification)with respect to the n=4(materials G, H, K and L)remembering that the C aromatic planes in the substrate. silicon carbide in the last sequence is the matrix itself SiC in the deposits was well crystallised, the size of the As discussed in a previous paper, the first interface has crystals being often limited by the thickness of the Sic a unique role in controlling the behaviour of the whole sublayer itself. The crystals are either of the cubic(3C)B interfacial sequence: fibre/Pyc, delle ng has to be modification or consisted of a sequence of disordered strong in order to allow multiple deflection at the dif polytypes, as shown in Fig. 5 and already reported by ferent other interfaces. Fibre surface bonding strength several authors.24. 25 The diffraction pattern(inset in Fig is related to the surface state of the fibre. In addition to 5)shows a straining of the reciprocal nodes along the the first interface(fibre/Py Ci) there exists two kinds of [lll] growth axis, i.e. perpendicular to the stacking interface characterised by a very different roughness fault plane henomena) in such multilayered(PyC/SiC)n inter The interface related to a Sic-deposit onto a Pyc c layer(e.g. PyCn-I/SiCn) was usually smooth, as already 00,2 mentioned, owing to the layered structure of pyr carbon and to its "covering capability"(tending to SIC t co0. 2 5c11 SiCAl 5n Fig. 5. Structure of the Sic-based sublayer. Inset: electron diffraction Fig 4. Growth of SiC on the surface of PyC layer: smoothness of the pattern centered on the contrasted crystal exhibiting one-dimensional
3.2. Nanostructure of the SiC-sublayers The SiC present in the multilayered (PyC/SiC)n interphase was always deposited onto pyrocarbon surfaces, roughly ¯at, as supported by a comparison of the pyrocarbon C(002) lattice fringes TEM images (recorded at about the same magni®cation) shown in Figs. 3 and 4. Fig. 4 shows that there is probably some relation between the orientation of the carbon aromatic planes in the substrate and that of the crystals in the SiC-deposit. Meanwhile, no precise relationship was found for the preferred growth directions of SiC (i.e. 111 for the cubic modi®cation and 00.1 for the hexagonal modi®cation) with respect to the C aromatic planes in the substrate. SiC in the deposits was well crystallised, the size of the crystals being often limited by the thickness of the SiC sublayer itself. The crystals are either of the cubic (3C) modi®cation or consisted of a sequence of disordered polytypes, as shown in Fig. 5 and already reported by several authors.24,25 The diraction pattern (inset in Fig. 5) shows a straining of the reciprocal nodes along the [111]c growth axis, i.e. perpendicular to the stacking fault plane. The ®rst SiC sublayer was observed to be often discontinuous, particularly when its thickness was low (i.e. 0.1 mm), as shown in Fig. 2. As a result, ``mechanical bridges'' were formed under such conditions, between the sublayers (inset in Fig. 2). Obviously, an optimising of the nucleation/growth processes should be further carried out in order to achieve thinner and smoother sublayers. 3.3. Multiple interfacing The number of interfaces in the (PyC/SiC)n multilayered interphase increases as n is raised: up to 8 when n=4 (materials G, H, K and L) remembering that the silicon carbide in the last sequence is the matrix itself. As discussed in a previous paper,11 the ®rst interface has a unique role in controlling the behaviour of the whole interfacial sequence: ®bre/PyC1 bonding has to be strong in order to allow multiple de¯ection at the different other interfaces.11 Fibre surface bonding strength is related to the surface state of the ®bre. In addition to the ®rst interface (®bre/PyC1) there exists two kinds of interface characterised by a very dierent roughness (important features regarding debonding and friction phenomena) in such multilayered (PyC/SiC)n interphases. The interface related to a SiC-deposit onto a PyClayer (e.g. PyCnÿ1/SiCn) was usually smooth, as already mentioned, owing to the layered structure of pyrocarbon and to its ``covering capability'' (tending to Fig. 4. Growth of SiC on the surface of PyC layer: smoothness of the interface. Fig. 5. Structure of the SiC-based sublayer. Inset: electron diraction pattern centered on the contrasted crystal exhibiting one-dimensional disordered polytypism. S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13 5
level the roughness of the substrate on which it has been first of all on the fibre surface bonding strength(fibre/ itself deposited). Apart the fibre/PyCl interface Py CI bonding strength). As received fibres are coated by obtained with the pristine fibre(which is the smoothest a thin oxide layer which produces a weak bonding. In interface) the second interface(PyC, SiC1) was system- contrast, with the treated fibres, the fibre surface bond atically that with the highest smoothness whatever the ing strength is raised. The detailed analysis of the nature of the material (since PyCl has been deposited on microcrack propagation paths fully supports the classi the fibre surface known to be smooth) fication of the materials in two families previously done The interface related to a pyrocarbon-deposit onto a on the basis of the mechanical behaviour and corre Sic-layer(e.g. SiCn-1/Py Cn)exhibited a high roughness sponding to the use of untreated or treated fibres, when the thickness of the Sic layer was large(owing to whatever the nature of the interphase (number and the tendency of Sic to grow as large-faceted crystals, as thickness of sublayers) hown in Fig. 5). It can be even discontinuous when the Sic layer is very thin. Under such conditions, there 3.4.1. Composites with untreated fibres occurs direct (and relatively strong) bonding between The microcrack paths exhibit two main features: (i)all the PyCn-I and PyCn layers(Fig. 2) the microcracks propagate up to the Nicalon fibre sur i) in the0°-fibr 3.4. Microcrack propagation path within a microscale there is no longer any bonding between the fibre and the multilayered interphase interphase, i.e. the fibre is debonded over its full length as shown in Fig. 7. These features occur whatever the As previously observed for single carbon interlayer, 5 nature of the interphase but as long as the fibres have matrix microcracks present in specimens loaded to fail- not been treated. Finally, for this first material family, ure(and especially crack-deflection-mechanisms)depend debonding and then sliding occurred mostly along the Fibre Matrix microcrack I Mode I Mode Il m Fig. 6. Multilayered interphase(material G with untreated Nicalon fibre) loaded to failure: (a) SEM micrograph on a polished longitudinal section showing the propagation path of a matrix microcrack (arrow)deflected in mode ll on fibre surface;(b) schematic showing the large residual crack opening and the final deflection at the fibre surface(open arrows indicate the loading direction)
level the roughness of the substrate on which it has been itself deposited). Apart the ®bre/PyC1 interface obtained with the pristine ®bre (which is the smoothest interface) the second interface (PyC1/SiC1) was systematically that with the highest smoothness whatever the nature of the material (since PyC1 has been deposited on the ®bre surface known to be smooth). The interface related to a pyrocarbon-deposit onto a SiC-layer (e.g. SiCnÿ1/PyCn) exhibited a high roughness when the thickness of the SiC layer was large (owing to the tendency of SiC to grow as large-faceted crystals, as shown in Fig. 5). It can be even discontinuous when the SiC layer is very thin. Under such conditions, there occurs direct (and relatively strong) bonding between the PyCnÿ1 and PyCn layers (Fig. 2). 3.4. Microcrack propagation path within a microscale multilayered interphase As previously observed for single carbon interlayer,15 matrix microcracks present in specimens loaded to failure (and especially crack-de¯ection-mechanisms) depend ®rst of all on the ®bre surface bonding strength (®bre/ PyC1 bonding strength). As received ®bres are coated by a thin oxide layer which produces a weak bonding. In contrast, with the treated ®bres, the ®bre surface bonding strength is raised. The detailed analysis of the microcrack propagation paths fully supports the classi- ®cation of the materials in two families previously done on the basis of the mechanical behaviour and corresponding to the use of untreated or treated ®bres, whatever the nature of the interphase (number and thickness of sublayers). 3.4.1. Composites with untreated ®bres The microcrack paths exhibit two main features: (i) all the microcracks propagate up to the Nicalon ®bre surface, as shown in Fig. 6 and (ii) in the 0-®bre-bundles, there is no longer any bonding between the ®bre and the interphase, i.e. the ®bre is debonded over its full length, as shown in Fig. 7. These features occur whatever the nature of the interphase but as long as the ®bres have not been treated. Finally, for this ®rst material family, debonding and then sliding occurred mostly along the Fig. 6. Multilayered interphase (material G with untreated Nicalon ®bre) loaded to failure: (a) SEM micrograph on a polished longitudinal section showing the propagation path of a matrix microcrack (arrow) de¯ected in mode II on ®bre surface; (b) schematic showing the large residual crack opening and the ®nal de¯ection at the ®bre surface (open arrows indicate the loading direction). 6 S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13
100 latrix PyCal 2 B(treated Nicalon fibre)after failure. TEM longitudinal Fig. 7. Material G(untreated Nicalon fibre). TEM cross-section(after a cohesive failure mode of the multilayered interpha failure) showing the fibre/multilayer debonding anywhere in the with mu ion (inset: schematic deflection in strong interface). fibres, which has been already described for single pyr- first interface which is very smooth, geometrically well carbon interphase. Owing to the high number of cracks defined and also very weak on a chemical view point 5 produced during this deflection, this damaging mode appears as a highly dissipating mechanism and results in a 3.4.2. Composites with treated fibres higher toughness. 2 It has been observed either for the Matrix microcrack deflections exhibit very different composite with a single thick carbon layer or for those features for the composites fabricated with the treated with(PyC/SiC)n multilayered interphases in which the fibres. First, it was never seen any matrix microcrack pyrocarbon sublayers are much thinner propagating up to the first interface, i.e. the fibre/PyC The second interface (PyCl/SiC1) owing Its interface. The fibre remains bonded over its full length smoothness was often observed to be the interface at (except very near the matrix microcracks) even at the (or near to) which the mode I/mode Il deflection occur- matrix crack saturation step; the fibre-matrix being red, as shown in Fig 9. The fibre is not debonded. Some- never uncoupled in a net manner(as observed in the times, microcracks were seen to be deflected at an other material family). As the fibre is being strained interface of much higher order. As an example, Fig. 10 under loading, debonding and then sliding (if there is shows the case of a deflection within the last carbon sub till any) no long occur along a well geometrically layer, far away from the fibre(favourable for protection defined surface but in a diffuse manner. Microcracks against oxidation) seem to burst into an infinity of nanometric-scale cracks Finally, a very common feature of the multilayer as as they are deflected in a pyrocarbon layer, parallel to deflector was the multideflection mode as seen for the fibre-axis. This was clearly identified to a shearing example in Fig. 8. The matrix microcrack underwent a failure mode in the case of a simple pyrocarbon inter- mode I/mode II deflection and then a mode Il/mode I layer 15 This is illustrated for material B in Fig. 8. and so on.. This multideflection mode was abundantly Deflecting within the whole thickness of the interphase seen in the case of the treated fibre. This is the tough (and not only as a single debonded surface) is a key ening-based mechanism suspected for that class of feature of composites fabricated with treated Nicalon material together with the matrix multiple cracking
®rst interface which is very smooth, geometrically well de®ned and also very weak on a chemical view point.15 3.4.2. Composites with treated ®bres Matrix microcrack de¯ections exhibit very dierent features for the composites fabricated with the treated ®bres. First, it was never seen any matrix microcrack propagating up to the ®rst interface, i.e. the ®bre/PyC1 interface. The ®bre remains bonded over its full length (except very near the matrix microcracks) even at the matrix crack saturation step; the ®bre-matrix being never uncoupled in a net manner (as observed in the other material family). As the ®bre is being strained under loading, debonding and then sliding (if there is still any) no long occur along a well geometrically de®ned surface but in a diuse manner. Microcracks seem to burst into an in®nity of nanometric-scale cracks as they are de¯ected in a pyrocarbon layer, parallel to the ®bre-axis. This was clearly identi®ed to a shearing failure mode in the case of a simple pyrocarbon interlayer.15 This is illustrated for material B in Fig. 8. De¯ecting within the whole thickness of the interphase (and not only as a single debonded surface) is a key feature of composites fabricated with treated Nicalon ®bres, which has been already described for single pyrocarbon interphase. Owing to the high number of cracks produced during this de¯ection, this damaging mode appears as a highly dissipating mechanism and results in a higher toughness.12 It has been observed either for the composite with a single thick carbon layer11 or for those with (PyC/SiC)n multilayered interphases in which the pyrocarbon sublayers are much thinner. The second interface (PyC1/SiC1) owing to its smoothness, was often observed to be the interface at (or near to) which the mode I/mode II de¯ection occurred, as shown in Fig. 9. The ®bre is not debonded. Sometimes, microcracks were seen to be de¯ected at an interface of much higher order. As an example, Fig. 10 shows the case of a de¯ection within the last carbon sublayer, far away from the ®bre (favourable for protection against oxidation). Finally, a very common feature of the multilayer as de¯ector was the multide¯ection mode as seen for example in Fig. 8. The matrix microcrack underwent a mode I/mode II de¯ection and then a mode II/mode I and so on...This multide¯ection mode was abundantly seen in the case of the treated ®bre. This is the toughening-based mechanism suspected for that class of material together with the matrix multiple cracking. Fig. 8. Material B (treated Nicalon ®bre) after failure. TEM longitudinal section exhibiting a cohesive failure mode of the multilayered interphase with multide¯ection (inset: schematic de¯ection in strong interface). Fig. 7. Material G (untreated Nicalon ®bre). TEM cross-section (after failure) showing the ®bre/multilayer debonding anywhere in the material. S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13 7
S. Bertrand et al. /Journal of the European Ceramic Society 20(2000)1-13 M 100nm mOnm Fig. 10. Material D(treated Nicalon fibre). TEM longitudinal section along the 0 bundle showing a mode I/mode Il deflection in the last a matrix microcrack at the second interface. Note that the first interface ( fibre/PyC1)is not debonded (TEM brightfield, sample loaded to failure) 4. Results-nanometer-scale multilayers processed by 4.1.2. The fibre surface and the first interface P-CVI The first material deposited on the fibre surface is systematically a pyrocarbon layer. Results of AES 4.1. Structure of the multilayered interphases ormed on the untreated and treated fibres, prior to any P-CVI infiltration, were per 4.1.1. Regularity and continuity of the layers formed in order to assess the chemical composition near Fig. 11(a)shows an example of a multilayered inter- the fibre surface. For the untreated fibres, a Hi-Nicalon phase processed at nanometric scale. This micrograph is yarn has been set in the reactor, and then heated under a cross section in a minicomposite obtained with as- vacuum, in order to reproduce the conditions experi- received Hi-Nicalon fibres and a multilayered inter- enced by the fibre prior to interphase deposition phase:(PyC3/SiC3o)10(3 and 30 being the PyC and Sic The surface of the untreated fibre is composed of a Si- layer thicknesses in nm and 10 the number of PyC/ Sic C-O mixture, 15 nm in thickness, assumed to be a sequences in the multilayer). A higher magnification SiO2+ free-C phase mixture. 20 The silica is formed dur [Fig. 11(b)] shows that pyrocarbon and Sic-based sub- ing the fibre heating under vacuum. When observed by layers are regular and continuous. The P-CVI process 4 TEM after the multilayer has been deposited, the fibre enables to deposit thin layers, parallel to the surface of surface looks rather smooth and clean. Fig. 12 is a high the fibres. It enables also to control the a- gaseous ratio resolution TEM micrograph of the first interface in a H2/MTS known to control the deposit composition and composite fabricated with untreated fibres and a(Pyc3/ crystallinity(here, the Sic-based layers do not consist of SiC3o)10 multilayer. The first pyrocarbon sublayer, but of a nanocrystalline SiC +C mixture). 4 composed of 7/8 carbon fringes, is lying directly on the Sharp interfaces between hard and compliant materials SiC nanocrystals forming the free surface of the fibre. are now accessible with that process. This has to be The oxide layer evidenced by AES is not observed after compared to those obtained at higher scale with I-CVI P-CVI, in the final material ig. 2). Layer flatness and thereafter interface sharpness The surface of the treated Hi-Nicalon fibre is composed, are suspected to be key features for obtaining a layered essentially, of a free-C layer, approximately 50 nm-thick. material with improved toughness Fig. 13 shows the interface of a material processed with a
4. ResultsÐnanometer-scale multilayers processed by P-CVI 4.1. Structure of the multilayered interphases 4.1.1. Regularity and continuity of the layers Fig. 11(a) shows an example of a multilayered interphase processed at nanometric scale. This micrograph is a cross section in a minicomposite obtained with asreceived Hi-Nicalon ®bres and a multilayered interphase: (PyC3/SiC30)10 (3 and 30 being the PyC and SiC layer thicknesses in nm and 10 the number of PyC/SiC sequences in the multilayer). A higher magni®cation [Fig. 11(b)] shows that pyrocarbon and SiC-based sublayers are regular and continuous. The P-CVI process14 enables to deposit thin layers, parallel to the surface of the ®bres. It enables also to control the -gaseous ratio H2/MTS known to control the deposit composition and crystallinity (here, the SiC-based layers do not consist of pure SiC but of a nanocrystalline SiC+C mixture).14 Sharp interfaces between hard and compliant materials are now accessible with that process. This has to be compared to those obtained at higher scale with I-CVI (Fig. 2). Layer ¯atness and thereafter interface sharpness are suspected to be key features for obtaining a layered material with improved toughness. 4.1.2. The ®bre surface and the ®rst interface The ®rst material deposited on the ®bre surface is systematically a pyrocarbon layer. Results of AES depth-pro®le analyses, performed on the untreated and treated ®bres, prior to any P-CVI in®ltration, were performed in order to assess the chemical composition near the ®bre surface. For the untreated ®bres, a Hi-Nicalon yarn has been set in the reactor, and then heated under vacuum, in order to reproduce the conditions experienced by the ®bre prior to interphase deposition. The surface of the untreated ®bre is composed of a SiC-O mixture, 15 nm in thickness, assumed to be a SiO2+free-C phase mixture.20 The silica is formed during the ®bre heating under vacuum. When observed by TEM after the multilayer has been deposited, the ®bre surface looks rather smooth and clean. Fig. 12 is a high resolution TEM micrograph of the ®rst interface in a composite fabricated with untreated ®bres and a (PyC3/ SiC30)10 multilayer. The ®rst pyrocarbon sublayer, composed of 7/8 carbon fringes, is lying directly on the SiC nanocrystals forming the free surface of the ®bre. The oxide layer evidenced by AES is not observed after P-CVI, in the ®nal material. The surface of the treated Hi-Nicalon ®bre is composed, essentially, of a free-C layer, approximately 50 nm-thick. Fig. 13 shows the interface of a material processed with a Fig. 10. Material D (treated Nicalon ®bre). TEM longitudinal section along the 0 bundle showing a mode I/mode II de¯ection in the last interfacial carbon sublayer. Fig. 9. Material L (treated Nicalon ®bre). Mode I/mode II de¯ection of a matrix microcrack at the second interface. Note that the ®rst interface (®bre/PyC1) is not debonded (TEM bright®eld, sample loaded to failure). 8 S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13
S. Bertrand et al. /Journal of the European Ceramic Society 20(2000)1-13 Fibre Interphase Matrix Fibre Fibre Matrix Fibre un 200nm Fig. 11. TEM ction of a minicomposite: as-received Hi-Nicalon/(PyC3/SiC3ohlo/SiC. Note the sharp interfaces controlled by the poor Sic crystallisation obtained by P-CVE (a) low magnification and (b) higher magnification on the interfacial sequence. reated Hi-Nicalon fibre. The residual carbon layer evi- denced by aEs is visible but with a thickness, usually thinner than 50 nm. It is in fact a bilayer. A rather dense carbon(less than 10 fringes)is lying directly on the surface of the fibre. Then a poorly organised carbon, at least 30 nm-thick, is observed. In Fig. 13, it can be seen than the first pyrocarbon sublayer (PyC1) is deposited on this poorly organised carbon with a more densely packed The two fibre surfaces are thus different, their rough ness meanwhile looking equivalent. In short, depending on the occurrence of a pretreatment or not, the first interface is either a SiC/PyC interface or a free-C/PyC 4.1.3. Nanostructure of the PyC sublayers The nanostructure of the Pyc-sublayers has been udied by TEM and optical microscopy. Values of the extinction angle(Ae),obtained for the PyC deposited by P-CVl, were all around 18-19. Each PyC sublayer grows onto surfaces made of well nanocrystallised SiC +C which exhibit some roughness at the nm-scale epted for the first sublayer). As shown in Figs. 12 13, the PyC first fills the concave parts of the Sic based substrate, at a distance the Pyc aromatic layers tend to deposit parallel to the mean surface of the Fig 12. Interface of an untreated hi-Nicalon fibre reinforced sic coated fibre and exhibit a pronounced anisotropy composite, with a(PyC3/SiC3o)o interphase
treated Hi-Nicalon ®bre. The residual carbon layer evidenced by AES is visible but with a thickness, usually, thinner than 50 nm. It is in fact a bilayer. A rather dense carbon (less than 10 fringes) is lying directly on the surface of the ®bre. Then a poorly organised carbon, at least 30 nm-thick, is observed. In Fig. 13, it can be seen than the ®rst pyrocarbon sublayer (PyC1) is deposited on this poorly organised carbon with a more densely packed stacking. The two ®bre surfaces are thus dierent, their roughness meanwhile looking equivalent. In short, depending on the occurrence of a pretreatment or not, the ®rst interface is either a SiC/PyC1 interface or a free-C/PyC1 interface. 4.1.3. Nanostructure of the PyC sublayers The nanostructure of the PyC-sublayers has been studied by TEM and optical microscopy. Values of the extinction angle (Ae),21 obtained for the PyC deposited by P-CVI, were all around 18±19. Each PyC sublayer grows onto surfaces made of well nanocrystallised SiC+C which exhibit some roughness at the nm-scale (excepted for the ®rst sublayer). As shown in Figs. 12 and 13, the PyC ®rst ®lls the concave parts of the SiCbased substrate, at a distance, the PyC aromatic layers tend to deposit parallel to the mean surface of the coated ®bre and exhibit a pronounced anisotropy. Fig. 11. TEM cross-section of a minicomposite: as-received Hi-Nicalon/(PyC3/SiC30)10/SiC. Note the sharp interfaces controlled by the poor SiC crystallisation obtained by P-CVI: (a) low magni®cation and (b) higher magni®cation on the interfacial sequence. Fig. 12. Interface of an untreated Hi-Nicalon ®bre reinforced SiC composite, with a (PyC3/SiC30)10 interphase. S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13 9
S. Bertrand et al. /Journal of the European Ceramic Society 20(2000)1-13 4.1.4. Nanostructure of the Sic-based sublayers fracture. The first interface remains meanwhile the most XRD analyses have been performed on planar important which controls the whole interphase beha deposits (i.e.(Py C/SiC)n multilayer deposited by P- viour. I Fig. 14 compares the brightfield and the carbon CVD). They show the coexistence of a and B Sic vari 002 dark field images of the same interphase. The latter ies as previously reported by Heurtevent 4 The is obtained as shown in the diffraction pattern, by measurement of the L111 from the spectra, is 18.8 nm. selecting the information carried by the 002 reflection by lectron probe mi cro-ana lysis means of an aperture. By this technique only the car- (EPMA), have shown that the Sic-based sublayers, are bon is in contrast. Despite their smooth aspect, this in fact, C+ SiC codeposits, as previously mentioned. 20 projection evidences the roughness of the sublayers. The In microcomposites, 4 the nanostructure of the Sic- multiple deflection is therefore expected to be limited based sublayers changes according to the thickness of the layer. Over 30 nm, a slight columnar microstructure 4.2. Crack propagation in the interphase apparent. The first Sic-based deposit (when e(sic) TEM low magnification images show clearly that a debonding occurs on most of the interfaces (arrows in Fig. 15). Even if the chemical interfacial bonding is strong, the strong fibre contraction when observed produces a high amount of debonding. As a Fibre Fig. 14. Interphase(Py Cxo/SiCso)o in a minicomposite with treated Fig. 13. Interfaces in treated Hi-Nicalon fibre-reinforced Sic compo- Hi-Nicalon fibres:(a) brightfield TEM image and(b)carbon 002-dark site. Inset is a 002 darkfield of the same interface showing three carbon field. Inset is the electron diffraction pattern of the multilayer with the layers at the treated Hi-Nicalon/(Py C2o/SiCso)o interface objective aperture position
4.1.4. Nanostructure of the SiC-based sublayers XRD analyses have been performed on planar deposits (i.e. (PyC/SiC)n multilayer deposited by PCVD). They show the coexistence of and SiC varieties as previously reported by Heurtevent.14 The measurement of the L111 from the spectra, is 18.8 nm. Microanalyses, by electron probe micro-analysis (EPMA), have shown that the SiC-based sublayers, are in fact, C+SiC codeposits, as previously mentioned.20 In microcomposites,14 the nanostructure of the SiCbased sublayers changes according to the thickness of the layer. Over 30 nm, a slight columnar microstructure is apparent. The ®rst SiC-based deposit (when e(SiC)<30 nm) is nanocrystallised, while, beyond 30 nm, a microcrystallised layer is superimposed on this ®rst amorphous layer. In minicomposites, the nanostructure of the SiCbased layers changes with the distance from the ®bre. The ®rst layers are nanocrystallised, while the layers, near the matrix, are microcrystallised. 4.1.5. The interfaces The multilayered structure is developed to multiply the number of interfaces in order to increase the work of fracture. The ®rst interface remains meanwhile the most important which controls the whole interphase behaviour.11 Fig. 14 compares the bright®eld and the carbon 002 dark ®eld images of the same interphase. The latter is obtained, as shown in the diraction pattern, by selecting the information carried by the 002 re¯ection by means of an aperture. By this technique, only the carbon is in contrast. Despite their smooth aspect, this projection evidences the roughness of the sublayers. The multiple de¯ection is therefore expected to be limited. 4.2. Crack propagation in the interphase These multilayered (PyC/SiC)n thin ®lms have been utilised as interphase in SiC/SiC composites. Uniaxial tension tests were performed at room temperature on the micro-, mini- and 2D-composites.14,20 After failure, the materials were studied by TEM and SEM to characterise the de¯ection mode and the interfacial behaviour. The matrix microcrack de¯ection mode depends on the treatment of the Hi-Nicalon ®bre. When the ®bre is not treated, a strong radial shrinkage occurs during the CVI processing at high temperature.25 TEM low magni®cation images show clearly that a debonding occurs on most of the interfaces (arrows in Fig. 15). Even if the chemical interfacial bonding is strong, the strong ®bre contraction when observed produces a high amount of debonding. As a Fig. 13. Interfaces in treated Hi-Nicalon ®bre-reinforced SiC composite. Inset is a 002 dark®eld of the same interface showing three carbon layers at the treated Hi-Nicalon/(PyC20/SiC50)10 interface. Fig. 14. Interphase (PyC20/SiC50)10 in a minicomposite with treated Hi-Nicalon ®bres: (a) bright®eld TEM image and (b) carbon 002-dark ®eld. Inset is the electron diraction pattern of the multilayer with the objective aperture position. 10 S. Bertrand et al. / Journal of the European Ceramic Society 20 (2000) 1±13