Journal of the European Ceramic Society 18(1 C 1998 Else Printed in Great Britain, All rig PII:s0955-2219(97)00I85-4 955221998s = Oxidizing environment Influence on the Mechanical Properties and microstructure of 2D-SIC/BN/ SiC Composites Processed by ICVI M. Leparoux, a L. Vandenbulcke, a V Serin, J. Sevely, S. Goujard and C. Robin-Brossec aCNRS-LCSR, Ic avenue de la Recherche Scientifique, F.-45071 Orleans, France CEMES-LOE. CNRS. 31055 Toulouse. france Societe Europeenne de Propulsion, Les Cinq Chemins-Le Haillan, F-33165 St-Medard-en-Jalles, france (Received 11 June 1997: accepted 22 October 1997) Abstract stability in severe environments. However it is now well established that the perfo 2D-SiC/Sic composites have been elaborated with vapor processed composites is strongly dependent BN interphases of diferent thicknesses which were on each component of the material and on the deposited on treated Nicalon( fibers by isothermal- processing parameters. 1, In a recent study on the isobaric chemical vapor infiltration from BCls isothermal/isobaric chemical vapor infiltration NHgH2 mixtures. Their mechanical behavior was (ICVD)of a boron nitride interphase, it was shown investigated at 600C in air and compared to the that the temperature, the gas flow and the inlet results obtained on similar SiC/C/SiC composites. composition conditions influenced both the BN Although these materials exhibit similar stress and organization and the nature of the interface with strain values at rupture when loaded at room tempera- the substratc. 3 Thcrcforc the knowledge and the ture, whatever the interphase, their thermomechanical control of each step of the composite manufactur- resistance depends on stress type, i.e. static or dynamic. ing are of prime importance Under static fatigue the BN interphases are more effi In order to obtain high mechanical properties cient than the pyrocarbon (PyC) ones. Thin BN inter- the composites generally require a compliant inter phases tend to maintain the interfacial properties. This phase between the fiber and the matrix. 4. 5 At first it result could be explained by the larger microcrack dis- was generally composed of a thin carbon film ances in the tows supporting the main part of the load, deposited on the fiber prior to the matrix infiltra according to a lower interfacial sliding resistance In tion0-8 This interphase exhibits a structure that contrast the materials with a PyC interphase, which allows sliding between the fiber and the matrix, and lave a much higher interfacial shear resistance at room then acts as a mechanical fuse. Nevertheless, the temperature, exhibit better thermomechanical beha carbon interphase is consumed dining envi under dynamic fatigue at 600C. The mechanical ronments, unless it is protected with a self-protec characteristics are related to the evolution of the fiber- tive matrix and/ or an external sealing coating. 9, 10 matrix interfacial zone which has been studied by When the whole composite is not protected,or SEM, TEM and EELS. @)1998 Elsevier Science when it is submitted to dynamic stresses that Limited. All rights reserved maintain the microcracks opening, the interpha is at least partially replaced by a glassy phase which strongly bonds the fiber to the matrix. 4, II 1 Introduction The result is a brittle behavior of the material In a previous study, the carbon interphase was High temperature structural applications, such as successfully rcplaccd by a boron nitride coating gas turbine engines or space plane thermal protection which was expected to be more oxidation resistant systems have generated a great interest in ceramic The BN interphase was elaborated by ICVI from matrix composites(CMC). These composites exhi- BCl3-NH3-H2 mixtures at a moderate temperature bit a combination of attractive physico-chemical of 700C. Thus, high strength and high strain properties such as high strength, toughness and have been obtained at room temperature with 715
Journal of the European Ceramic Society 18 (1998) 115-723 6 1998 Elsevier Science Limited PII: SO955-2219(97)00185-4 Printed in Great Britain. All rights reserved 0955-2219/98/$19.00 + 0.00 Oxidizing Environment Influence on the Mechanical Properties and Microstructure of 2D-SiC/BN/SiC Composites Processed by ICVI M. Leparoux,a L. Vandenbulcke,” V. Seriqh J. Sevely,h S. Goujard” and C. Robin-Brosse” “CNRS-LCSR, 1C avenue de la Recherche Scientifique, F-45071 Orlkans, France ‘CEMES-LOE, CNRS, 3 1055 Toulouse, France “SociCtC Europkenne de Propulsion, Les Cinq Chemins-Le Haillan, F-33 165 St-MCdard-en-Jalles, France (Received 11 June 1997; accepted 22 October 1997) Abstract ZD-Sic/Sic composites have been elaborated with BN interphases of d@erent thicknesses which were deposited on treated ‘Nicalon@‘bers by isothermalisobaric chemical vapor infiltration from BClr NH,H2 mixtures. Their mechanical behavior was investigated at 600°C in air and compared to the results obtained on similar SiCjCjSiC composites. Although these materials exhibit similar stress and strain values at rupture when loaded at room temperature, whatever the interphase, their thermomechanical resistance depends on stress type, i.e. static or dynamic. Under static fatigue the BN interphases are more eficient than the pyrocarbon (PyC) ones. Thin BN interphases tend to maintain the interfacial properties. This result could be explained by the larger microcrack distances in the tows supporting the main part of the load, according to a lower interfacial sliding resistance. In contrast the materials with a PyC interphase, which have a much higher interfacial shear resistance at room temperature, exhibit better thermomechanical behavior under dynamic fatigue at 600°C. The mechanical characteristics are related to the evolution of the$bermatrix interfacial zone which has been studied by SEM, TEM and EELS. 0 1998 Elsevier Science Limited. All rights reserved 1 Introduction High temperature structural applications, such as gas turbine engines or space plane thermal protection systems have generated a great interest in ceramic matrix composites (CMC). These composites exhibit a combination of attractive physico-chemical properties such as high strength, toughness and 715 stability in severe environments. However it is now well established that the performance of chemical vapor processed composites is strongly dependent on each component of the material and on the processing parameters. l,* In a recent study on the isothermal/isobaric chemical vapor infiltration (ICVI) of a boron nitride interphase, it was shown that the temperature, the gas flow and the inlet composition conditions influenced both the BN organization and the nature of the interface with the substrate.3 Therefore the knowledge and the control of each step of the composite manufacturing are of prime importance. In order to obtain high mechanical properties, the composites generally require a compliant interphase between the fiber and the matrix.4,5 At first it was generally composed of a thin carbon film deposited on the fiber prior to the matrix infiltration.“* This interphase exhibits a structure that allows sliding between the fiber and the matrix, and then acts as a mechanical fuse. Nevertheless, the carbon interphase is consumed in oxidizing environments, unless it is protected with a self-protective matrix and/or an external sealing coating.9,‘0 When the whole composite is not protected, or when it is submitted to dynamic stresses that maintain the microcracks opening, the interphase is at least partially replaced by a glassy phase which strongly bonds the fiber to the matrix.4,” The result is a brittle behavior of the material. In a previous study, ‘* the carbon interphase was successfully replaced by a boron nitride coating which was expected to be more oxidation resistant. The BN interphase was elaborated by ICVI from BCls-NH3-H2 mixtures at a moderate temperature of 700°C. Thus, high strength and high strain have been obtained at room temperature with
2D-SiC/SiC composites in which the fibers were The boron nitride interphase was infiltrated hemically treated prior to any deposition stage the within the porous fiber preforms by the icvi pro surface of these treated Nicalon fibers exhibited a cess. BCl3-NH3-H2 gas mixtures were used at a different composition over significant thickness. 3 moderate temperature of 700 C and under reduced However it has been shown that the interfacial slid pressure according to a procedure which has been g resistance increased with the interphase thick- widely described elsewhere. 3, 16 Various interphase ness at room temperature, even if a bn thickness thicknesses, in the range 0.2-0.7 um, have been variation in the 0-2-0-7 um range did not induce deposited. The BN-coated Nicalon@ preforms any differences in the ultimate strength and strain. were further densified with Sic by a classical ICVI Thus crack saturation occurred for the lowest process from methyltrichlorosilane-H2 mixtures interphase thickness, and the microcracks spacing Moreover, after the specimen manufacturing, a sic in the tensile stress direction(within the 0 fibers- coating was deposited to prevent any direct initial tows)decreased as the bn thickness increased. 12 oxygen access to the interphase and interfaces The first objective of this study was to study the Some mechanical characteristics of SiC/Sic hemical and structural evolution of the interphase composites that include a pyrocarbon(PyC) inter and interfaces of 2D-SiC/BN SiC composite when phase deposited on treated Nicalon fibers are exposed to mechanical stress in an oxidizing envi- provided for comparison. In all cases, the materials ronment. It is indeed well established that the fiber/ were overcoated by SiC after machining matrix interface mainly governs the thermo- mechanical behavior of such composites. 7, 14, 5 2.2 Mechanical tests Therefore, various mechanical tests have been per- The mechanical behavior of the 2D composites was formed such as static bending and tensile tests, as evaluated at 600C in air by static bending and well as dynamic fatigue in tension-tension. The tensile tests as well as dynamic tension-tension nfluence of the interphase thickness on the fatigue tests. For these tests, bars were machined mechanical behavior was also examined. These with their main direction parallel to one of the fiber mechanical characterizations were carried out in orientation. The geometry and dimensions of the air at a temperature of 600C which corresponds to tensile and bending specimens are presented in a moderate temperature for the occurrence of hex- Fig. 1. The static fatigue tests were performed on agonal-BN oxidation. After failure of the compo- 4-points bending machine prototype having major site, the interphase and interfaces were examined at and minor spans of 5-08 and 2.54 cm. The tensile different scales using scanning electron microScopy fatigue behavior was studied with a hydraulic but also transmission electre ppy for the machine(Instron Dynamic 8501) equipped with microstructure and energy loss spectroscopy for hydraulic grips cooled with water. These charac- the chemistry of each interfacial zone. In addition, terizations were performed within a furnace (aET)in the mechanical results obtained for the SiC/BN/ ambient air. The tensile strain was measured with an compared to the results obtained on similar mat- length. For the dynamic tests in tension-tension, the erials with a pyrocarbon interphase frequency was adjusted at the desired value, 2 or 20 Hz in the present case, after few cycles at 0.02 Hz. 2 Experimental procedures 2.3 Microstructure and chemistry of the interfacial zones 2.1 Materials The interfacial microstructures were studied by The 2D-SiC/BN/SiC composites were prepared as transmission electron microscopy (TEm)whereas rectangular plates(160×80×3mm3)from2Dpre forms consisting of stacks of fabrics made with pressed together with graphite tooling to obtain a constant nominal fiber content of about 37 vol% The Nicalon fibers(NLM 202)were previousl chemically treated (proprietary treatment) by the treated fiber presented a different composition over 9+ a significant thickness and the surface was free carbon rich. 3 At room temperature it has been checked that this treatment did not decrease the Fig. 1. Geometry and dimensions, in mm, of the(a)tensile mechanical properties of the fibers (b )bending sp
716 M. Leparoux et al. 2D-Sic/Sic composites in which the fibers were The boron nitride interphase was infiltrated chemically treated prior to any deposition stage. The within the porous fiber preforms by the ICVI prosurface of these treated Nicalonm fibers exhibited a cess. BC13-NH3-H2 gas mixtures were used at a different composition over significant thickness.13 moderate temperature of 700°C and under reduced However it has been shown that the interfacial slid- pressure according to a procedure which has been ing resistance increased with the interphase thick- widely described elsewhere.3a16 Various interphase ness at room temperature, even if a BN thickness thicknesses, in the range 0.2-0.7 pm, have been variation in the 0.24.7pm range did not induce deposited. The BN-coated Nicalon@ preforms any differences in the ultimate strength and strain. were further densified with SIC by a classical ICVI Thus crack saturation occurred for the lowest process from methyltrichlorosilane-H2 mixtures. interphase thickness, and the microcracks spacing Moreover, after the specimen manufacturing, a SIC in the tensile stress direction (within the 0” fibers- coating was deposited to prevent any direct initial tows) decreased as the BN thickness increased.12 oxygen access to the interphase and interfaces. The first objective of this study was to study the chemical and structural evolution of the interphase and interfaces of 2D-SiC/BN/SiC composite when exposed to mechanical stress in an oxidizing environment. It is indeed well established that the fiber/ matrix interface mainly governs the thermomechanical behavior of such composites.7,14,‘5 Therefore, various mechanical tests have been performed such as static bending and tensile tests, as well as dynamic fatigue in tension-tension. The influence of the interphase thickness on the mechanical behavior was also examined. These mechanical characterizations were carried out in air at a temperature of 600°C which corresponds to a moderate temperature for the occurrence of hexagonal-BN oxidation. After failure of the composite, the interphase and interfaces were examined at different scales using scanning electron microscopy but also transmission electron microscopy for the microstructure and energy loss spectroscopy for the chemistry of each interfacial zone. In addition, the mechanical results obtained for the SiC/BN/ SIC composites with various BN thicknesses were compared to the results obtained on similar materials with a pyrocarbon interphase. Some mechanical characteristics of Sic/Sic composites that include a pyrocarbon (PyC) interphase deposited on treated Nicalon@ fibers are provided for comparison. In all cases, the materials were overcoated by SIC after machining. 2.2 Mechanical tests The mechanical behavior of the 2D composites was evaluated at 600°C in air by static bending and tensile tests as well as dynamic tension-tension fatigue tests. For these tests, bars were machined with their main direction parallel to one of the fiber orientation. The geometry and dimensions of the tensile and bending specimens are presented in Fig. 1. The static fatigue tests were performed on a 4-points bending machine prototype having major and minor spans of 5.08 and 2.54cm. The tensile fatigue behavior was studied with a hydraulic machine (Instriin Dynamic 8501) equipped with hydraulic grips cooled with water. These characterizations were performed within a furnace (AET) in ambient air. The tensile strain was measured with an extensometer (2620-603 Instron) with a 25 mm gauge length. For the dynamic tests in tension-tension, the frequency was adjusted at the desired value, 2 or 20 Hz in the present case, after few cycles at 0.02 Hz. 2 Experimental procedures 2.1 Materials The 2D-SiC/BN/SiC composites were prepared as rectangular plates (160x 80x3 mm3) from 2D preforms consisting of stacks of fabrics made with Nicalon@ fibers. These stacks were maintained and pressed together with graphite tooling to obtain a constant nominal fiber content of about 37 ~01%. The Nicalon@ fibers (NLM 202) were previously chemically treated (proprietary treatment) by the Sociite Europeenne de Propulsion (SEP). The treated fiber presented a different composition over a significant thickness and the surface was free carbon rich.13 At room temperature, it has been checked that this treatment did not decrease the mechanical properties of the fibers. 2.3 Microstructure and chemistry of the interfacial zones The interfacial microstructures were studied by transmission electron microscopy (TEM) whereas 20 I- I- (4 3 60:s 60 P 60 (b) Fig. 1. Geometry and dimensions, in mm, of the (a) tensile and (b) bending specimens
Oxidizing environment influence on 2D-SiC/BN SiC composites microchemical analyses were achieved by electron energy loss spectrometry (EFL S). These experi- ments were performed on a Philips CM30ST 250● microscope fitted with a Gatan 666 spectrometer carbon EELS analyses have been generally carried out F with a probe size about 20 nm in diameter, reduced to 7 nm for the study of thin sublayers. The energy resolution of the spectrometer was about 1. 5 cV. The samples were taken of a tensile- tested speci men, near the failure surface. Thin foils (S500A were prepared in a conventional way by mechan ical polishing followed by Ar ion milling Lifetime(h) 3 Results Fig. 2. Static fatigue behavior in flexion for 2D-SiC/SIC 3.1 Mechanical results composites with various interphases deposited on treated All the mechanical tests were performed in the fractured when the test was stopped; the pyrocarbon thickness non-linear domain of the composite at 600oC in was in the 0. 1-02 um range) air. At this temperature, a crack sealing by oxida- tion of the SiC matrix is unexpected. 7, 18 On the other hand, at such low temperature no decrease of frequency was adjusted as a function of the applied the mechanical properties of the treated Nicalon@ stress. The results obtained for similar composites fiber can occur due to its thermal decomposi- with a pyrocarbon interphase are also summarized tion. 18, 19 In that way, the results are considered in Table 1. It can be noted that various types of to be mainly representative of the evolution of SiC/C/SiC composites, corresponding to different the fibcr/matrix interfacial zonc in an oxidizing processing conditions, have been used for the environment characterization at 120 MPa. they broke at differ ent time in the 15 to 24 h range. For comparison,a 3.. I Static bending tests composite with a 0.7 um bn thickness did not The 2D-SiC/BN/SiC composites were made with break after about 1 6x 10 cycles at 20 Hz (22 h) treated Nicalon fibers and various interphase when loaded between 0 and 150 MPa, but at room hicknesses of 0.7, 0-4 and 0.2 um. They were sub- temperature mitted to a static flexion fatigue testing in air at Even though only a few dynamic tests have 600C. Two specimens of each composites, includ- performed and only one specimen has been ing different thicknesses as well as a pyr- acterized at each stress level, it appears that tI ocarbon interphase, were tested using a constant SiC/BN SiC materials break before the Sic/c/Sic applied load to determine the time to cause frac- composites. The result at 200 MPa may be not sig ture as a function of the applied stress. the values nificant because of the short duration of the test for the first specimen to fail are plotted in Fig. 2. obviously, the composites retain their mechanical Both samples of the composite made with a 0.2 strength for a longer duration as the stress decrea- BN interphase were not fractured after more ses. On the other hand, results for SiC/BN/SiC 300h at 150 MPa. It must be noted that the ulti- would suggest that lower bn thicknesses lead to mate tensile strength at room temperature was better mechanical results at 600C in an oxidizing about 300-320 MPa for all materials environment. however this behavior needs to be The static fatigue tests in air at 600C show that confirmed the oxidation resistance of the Sic/Sic composites Furthermore, static fatigue tests in tension have is greatly improved when a BN interphase is used bcen performed on a SiC/BN SiC composite. In instead of the pyrocarbon one. Moreover the fail- that case, the interphase thickness was 0.7 um and re of the site appears later when the only one specimen could be loaded at different applied stress and the interphase thickness are low. stress levels. The results are reported in Table 2. When compared with flexion tests (Fig. 2), they 3.1.2 Tensile fatigue characterizations confirm a higher sensitivity to fatigue tests in ten One specimen of each material was tested in ten- sion than in flexion, as currently observed. If we sion-tension at different stress levels for the compare with the results of Table 1 obtaine dynamic characterizations. The results are reported tension-tension, the results of the static as minimum time to cause failure(Table 1). The reported in Table 2 are obviously far better
Oxidizing environment influence on 2D-SiC/BN/SiC composites 717 microchemical analyses were achieved by electron energy loss spectrometry (EELS). These experiments were performed on a Philips CM30ST microscope fitted with a Gatan 666 spectrometer. EELS analyses have been generally carried out with a probe size about 20 nm in diameter, reduced to 7 nm for the study of thin sublayers. The energy resolution of the spectrometer was about 1.5 eV. The samples were taken of a tensile-tested specimen, near the failure surface. Thin foils (ISOOA) were prepared in a conventional way by mechanical polishing followed by Ar+ ion milling. 3 Results 3.1 Mechanical results All the mechanical tests were performed in the non-linear domain of the composite at 600°C in air. At this temperature, a crack sealing by oxidation of the SIC matrix is unexpected.i7J8 On the other hand, at such low temperature no decrease of the mechanical properties of the treated Nicalon@ fiber can occur due to its thermal decomposition.18J9 In that way, the results are considered to be mainly representative of the evolution of the fiber/matrix interfacial zone in an oxidizing environment. 3.1.1 Static bending tests The 2D-SiC/BN/SiC composites were made with treated Nicalon@ fibers and various interphase thicknesses of 0.7, 0.4 and 0.2,~rn. They were submitted to a static flexion fatigue testing in air at 600°C. Two specimens of each composites, including different BN thicknesses as well as a pyrocarbon interphase, were tested using a constant applied load to determine the time to cause fracture as a function of the applied stress. The values for the first specimen to fail are plotted in Fig. 2. Both samples of the composite made with a O-2 ,um BN interphase were not fractured after more than 300 h at 150 MPa. It must be noted that the ultimate tensile strength at room temperature was about 300-320 MPa for all materials. The static fatigue tests in air at 600°C show that the oxidation resistance of the Sic/Sic composites is greatly improved when a BN interphase is used instead of the pyrocarbon one. Moreover the failure of the composite appears later when the applied stress and the interphase thickness are low. 3.1.2 Tensile fatigue characterizations One specimen of each material was tested in tension-tension at different stress levels for the dynamic characterizations. The results are reported as minimum time to cause failure (Table 1). The 250 - 0 is % 3 200 - 0 g v) * BN=0.4pm * A A BN=0.2pm A 150 - 0 * A+ #aI I a 1481, 10 2 6 100 2 6 1000 Lifetime (h) Fig. 2. Static fatigue behavior in flexion for 2D-SiC/SiC composites with various interphases deposited on treated Nicalon@ fibers (the arrow indicates that the specimen was not fractured when the test was stopped; the pyrocarbon thickness was in the 0.14.2 pm range). frequency was adjusted as a function of the applied stress. The results obtained for similar composites with a pyrocarbon interphase are also summarized in Table 1. It can be noted that various types of SiCjCjSiC composites, corresponding to different processing conditions, have been used for the characterization at 120 MPa. They broke at different time in the 15 to 24 h range. For comparison, a composite with a 0.7pm BN thickness did not break after about 1.6x lo6 cycles at 20 Hz (22 h) when loaded between 0 and 150 MPa, but at room temperature. Even though only a few dynamic tests have been performed and only one specimen has been characterized at each stress level, it appears that the SiC/BN/SiC materials break before the Sic/C/Sic composites. The result at 200MPa may be not significant because of the short duration of the test. Obviously, the composites retain their mechanical strength for a longer duration as the stress decreases. On the other hand, results for SiC/BN/SiC would suggest that lower BN thicknesses lead to better mechanical results at 600°C in an oxidizing environment. However, this behavior needs to be confirmed. Furthermore, static fatigue tests in tension have been performed on a SiC/BN/SiC composite. In that case, the interphase thickness was 0.7,~rn and only one specimen could be loaded at different stress levels. The results are reported in Table 2. When compared with flexion tests (Fig. 2) they confirm a higher sensitivity to fatigue tests in tension than in flexion, as currently observed. If we compare with the results of Table 1 obtained in tension-tension, the results of the static tests reported in Table 2 are obviously far better
718 M. Leparoux et al Table 1. Minimum number of cycles and duration to cause failure of 2 D-SiC/SiC composites with different interphases. They were tested in tension-tension, in air at 600 C, with various applied stresses bN thickness PyC MPa) 01-02m μm 02 4000(7h30min)68400(9h30min)108000-172800(1524h) 9000(h15min) 19800(2h45min) 0-200 2000(18min) 9500(8min) Table 2. Time to cause failure of 2D-SiC/ BN/SiC composite microscopy (TEM)and clectron energy loss spec- submitted to static tensile stress. The interphase thickness was troscopy(EELS). The analyses were radially per cates no rupture formed, from the fiber to the matrix through the BN interphase and its interfaces both with the bn thickne (MPa) treated Nicalon fiber and the Sic matrix ccording to a procedure which has been described elsewhere. 5 It is interesting to recall that the 145 15h I h interfaces were initially carbon-rich, before any exposure to the oxidative environment. 5Because 3.2 Evolution of the fiber/matrix interfacial zones our interest was principally focused on the chem in an oxidizing environment istry of the interfacial zone, no attempt was made to characterize the matrix and the fiber 3.2.1 SEM observations 3.2.2.1 The bn interphase. The thickness of These observations have been performed on a the Bn interphase is homogeneous(about 0.5 um) sample made with a bn thickness of 0.5 um which and no voids are observed(Fig. 9). The eels had failed after a tensile test of about 70 h at 600 C. spectra of this region show boron and nitrogen In that case another test procedure was employed edges(Fig. 10) typical of these elements involved that permits to evaluate quickly the performances of in hexagonal boron nitride. u In addition, they a new material: the applied stress was increased by evidence significant amounts of carbon and oxy increments of 10 MPa. After each step the following gen in the bn phase. Silicon was not dctcctcd, increment was applied when the strain of the com- however it could be below the detection limit of posite nearly reached a stable state. Finally, failure of EELS (<I at%). The concentration of B, C, the composite occurred at 170 MPa. and O was determined by a classical quantitative SEM micrographs of polished cross-sections analysis technique. The average atomic com show that the bn coating is uniformly deposited position was 40% N, 48%B, 7%C and 5%O inside the preform [Fig 3(aH(c). The fiber/matrix In comparison with the average composition interfacial zone either appears smooth or presents measured on the same composite tested at room some crystalline or amorphous masses [Fig 4(a) temperature (41. 5% N, 39%B, 10% C and and(b)]. Microcracks are observed at the fiber/ Bn 9.5%O), 5 we note a decrease of the o, C and N interface but also in the interphase itself and at the contents BN/matrix interface Fig. 5(a) and(b). They 3. 2.2.2 The fiber/bn interface. The TEM propagate randomly in these three regions, on the micrographs show many large decohesion at this contrary to the results obtained at room tempera- interface as clearly shown in Fig. 9. Otherwise, ture where these cracks mainly propagate at the when the fiber and the interphase are still bonded, fiber/BN interface. 12 Many links that present aa bright zone of about 10 nm thick borders the vitreous appearance are seen between Bn and the fiber. This sublayer is mainly composed of silicon fiber(Fig. 6)and between BN and the matrix. These and oxygen with traces of boron and carbon. The particular glassy phases are also observed in cracks near-edge fine structure of the Si-L2,3 edge(double that propagate into the matrix between two neigh- peak at 109 and 116 ev)(Fig. 11)and the dis boring fibers(Fig. 7). The fracture surface shows symmetrical plasmon peak at 22. 5 eV in the range gencrally a short pull-out lcngth [Fig 8(a)and(b)]. of low cncrgy loss(Fig. ll, inset) are consistent Moreover, it appears that the Bn coating adheres with the presence of silica at this interface. The both to the fiber and the matrix [Figs 5(b)and 8(c)]. quantitative analyses give an avcrage Si/ o atomic ratio of 0. 53. It can be recalled here th 3. 2. 2 Microstructure and chemistry of the fiber/ treated Nicalon@ NLM 202 fiber present matrix interfacial zone carbon-rich surface which was maintained at the Further microstructure investigations have been fiber/BN interface, at the end of the composite made on the same sample by transmission electron manufacturing. I5
718 M. Leparoux et al. Table 1. Minimum number of cycles and duration to cause failure of 2D-Sic/Sic composites with different interphases. They were tested in tension-tension, in air at 6OO”C, with various applied stresses Stress (MPa) Frequency (Hz) 0.7 pm BN thickness 0.4 pm 0.2 pm PYC 0.14.2 pm Cl20 G150 &200 2 54 000 (7 h 30 min) 68 400 (9 h 30 min) 108 000-172 800 (15-24 h) 2 9000 (lh 15 min) 19 800 (2 h 45 min) 20 22 000 (18min) 9500 (8 min) Table 2. Time to cause failure of 2D-SiC/BN/SiC composites submitted to static tensile stress. The interphase thickness was 0.7pm. R indicates a rupture of the materials and NR indicates no rupture Stress BN thickness (MPa) 0.7 pm 120 > 1OOh 145 15h 185 lh NR R R 3.2 Evolution of the fiber/matrix interfacial zones in an oxidizing environment 3.2.1 SEA4 observations These observations have been performed on a sample made with a BN thickness of 0.5pm which had failed after a tensile test of about 70 h at 600°C. In that case another test procedure was employed that permits to evaluate quickly the performances of a new material: the applied stress was increased by increments of 10 MPa. After each step the following increment was applied when the strain of the composite nearly reached a stable state. Finally, failure of the composite occurred at 170 MPa. SEM micrographs of polished cross-sections show that the BN coating is uniformly deposited inside the preform [Fig. 3(a)-(c)]. The fiber/matrix interfacial zone either appears smooth or presents some crystalline or amorphous masses [Fig. 4(a) and (b)]. Microcracks are observed at the fiber/BN interface but also in the interphase itself and at the BN/matrix interface [Fig. 5(a) and (b)]. They propagate randomly in these three regions, on the contrary to the results obtained at room temperature where these cracks mainly propagate at the fiber/BN interface. l2 Many links that present a vitreous appearance are seen between BN and the fiber (Fig. 6) and between BN and the matrix. These particular glassy phases are also observed in cracks that propagate into the matrix between two neighboring fibers (Fig. 7). The fracture surface shows generally a short pull-out length [Fig. 8(a) and (b)]. Moreover, it appears that the BN coating adheres both to the fiber and the matrix [Figs 5(b) and 8(c)]. 3.2.2 Microstructure and chemistry of the$ber/ matrix interfacial zone Further microstructure investigations have been made on the same sample by transmission electron microscopy (TEM) and electron energy loss spectroscopy (EELS). The analyses were radially performed, from the fiber to the matrix through the BN interphase and its interfaces both with the treated Nicalon@ fiber and the SIC matrix, according to a procedure which has been described elsewhere.15 It is interesting to recall that the interfaces were initially carbon-rich, before any exposure to the oxidative environment.15 Because our interest was principally focused on the chemistry of the interfacial zone, no attempt was made to characterize the matrix and the fiber. 3.2.2.1 The BN interphase. The thickness of the BN interphase is homogeneous (about 0.5 pm) and no voids are observed (Fig. 9). The EELS spectra of this region show boron and nitrogen edges (Fig. 10) typical of these elements involved in hexagonal boron nitride.20 In addition, they evidence significant amounts of carbon and oxygen in the BN phase. Silicon was not detected, however it could be below the detection limit of EELS (< 1 at%). The concentration of B, C, N and 0 was determined by a classical quantitative analysis technique. 21 The average atomic composition was 40% N, 48% B, 7% C and 5% 0. In comparison with the average composition measured on the same composite tested at room temperature (41.5% N, 39% B, 10% C and 9.5% 0),15 we note a decrease of the 0, C and N contents. 3.2.2.2 The fiber/BN interface. The TEM micrographs show many large decohesions at this interface as clearly shown in Fig. 9. Otherwise, when the fiber and the interphase are still bonded, a bright zone of about 10nm thick borders the fiber. This sublayer is mainly composed of silicon and oxygen with traces of boron and carbon. The near-edge fine structure of the Si-L2,3 edge (double peak at 109 and 116 eV) (Fig. 11) and the dissymmetrical plasmon peak at 22.5 eV in the range of low energy loss (Fig. 11, inset) are consistent with the presence of silica at this interface. The quantitative analyses give an average SijO atomic ratio of 0.53. It can be recalled here that the treated Nicalon@ NLM 202 fiber presented a carbon-rich surface which was maintained at the fiber/BN interface, at the end of the composite manufacturing.15
Oxidizing environment influence on 2D-SiC/ BN/ Sic composite 719 3.2.2.3 The BN/Sic-matrix interface. Figure 9 the near-edge fine structure of the Si-L2, 3 edge may shows the interface between the SiC matrix and the suggest that Si atoms are involved in both Sic and BN interphase. It is coarse and porous with a SiO2 5(Fig. 12). An additional hrem study could thickness of about 35 nm. Although the feature of ermit to afford more information about the rela this zone appeared similar to the interface observed tive organization of SiC, SiO2 and C. Furthermore on the same composite tested at room temperature, close to this intcrface, inside the boron nitride the composition is significantly modified. EELs some silica has also been detected analyses have revealed high concentration of oxy- 3. 2.2. 4 Analysis of the cracks. Along a decohe gen (50 at%), silicon (34 at%)and carbon(16 sion at the fiber/BN interface, EELS analyses at%), instead of carbon in the original compo revealed the presence of Si, O and B atoms. Car site.5 Additionally, the damped two pre-peaks in bon was also revealed, but it remains under the quantitative limit, while no nitrogen was detected The average atomic ratio O/Si was equal to 1 3 and the boron concentration was below 5 at% On the other hand the chemical characterization of a crack that phase allows to determine an average atomic com position of 34% B, 28%N, 20%O, 10%C and 8% Si, The near edge structure of the Si-L2 3 peak exhibits a double peak on the edge onset. It is similar to the one previously seen at the fiber/BN interface and corresponds to the presence, in this crack, of Si-o bonds in a SiO, structure 50 Although, in both cases, the presence of boron oxide could be considered from the surstoichiometry 2 10 um 1 10 um Fig3.(a) Cross-section of a 2D-SiC/BN/SiC composite,(b) Fig. 4. Interfacial zone showing crystalline or amorphous in the external part and (c)in the middle of a fibers tow
Oxidizing environment influence on 2D-SiC/BN/SiC composites 719 3.2.2.3 The BN/SiC-matrix interface. Figure 9 shows the interface between the SIC matrix and the BN interphase. It is coarse and porous with a thickness of about 35nm. Although the feature of this zone appeared similar to the interface observed on the same composite tested at room temperature, the composition is significantly modified. EELS analyses have revealed high concentration of oxygen (50 at%), silicon (34 at%) and carbon (16 at%), instead of carbon in the original composite.i5 Additionally, the damped two pre-peaks in (b) the near-edge fine structure of the Si-L2,3 edge may suggest that Si atoms are involved in both Sic and Si0215 (Fig. 12). An additional HREM study could permit to afford more information about the relative organization of Sic, Si02 and C. Furthermore, close to this interface, inside the boron nitride, some silica has also been detected. 3.2.2.4 Analysis of the cracks. Along a decohesion at the fiber/BN interface, EELS analyses revealed the presence of Si, 0 and B atoms. Carbon was also revealed, but it remains under the quantitative limit, while no nitrogen was detected. The average atomic ratio O/Si was equal to 1.3 and the boron concentration was below 5 at%. On the other hand, the chemical characterization of a crack that propagates radially into the interphase allows to determine an average atomic composition of 34% B, 28% N, 20% 0, 10% C and 8% Si. The near edge structure of the Si-L2,3 peak exhibits a double peak on the edge onset. It is similar to the one previously seen at the fiber/BN interface and corresponds to the presence, in this crack, of Si-0 bonds in a Si02 structure.** Although, in both cases, the presence of boron oxide could be considered from the surstoichiometry (4 (b) Fig. 3. (a) Cross-section of a 2D-SiC/BN/SiC composite, (b) Fig. 4. Interfacial zone showing crystalline or amorphous in the external part and (c) in the middle of a fibers tow. masses after a tensile test of 70 h, at 600°C in air
M. Leparoux et al. in boron as compared with nitrogen, that cannot the volatility of boron oxide, it is not surprising to be evidenced by our analyses find such a phase in different parts of the compo- site, especially in the tows, inside cracks perpend cular or parallel to the fibers (Figs 6 and 7) 4 Discussion Sometimes it can even link the cracks walls. The EELS analysis along a decohesion at the fiber /BN The application of a compliant interphase such as interface confirms that the Si, O and B elements are carbon or boron nitride on treated Nicalon M fibers present(Section 3. 2. 2, Analysis of the cracks has been found to improve the mechanical beha- analysis of a crack inside the bn interphase vior of the SiC/SiC composites at room tempera- also an incrcasc in oxygen contcnt and the ture. The drawbacks of carbon are obvious in presence with Si-o bonds formation oxidizing atmosphere, and bn should exhibil a From EELS study, it appears that the bulk better oxidation resistance, especially at moderate boron nitride interphase can be modified at 600C temperature such as 600C. However the ICVI-BN in air. However the decrease of the o, C and N deposits always present a turbostratic structure. contents should be examined by further analyses, Even if the temperature of 600 C is low for the particularly at different distances from the far end onset of hexagonal-BN oxidation, poorer lattice of the cracks, i.e. different conditions of ambient organization might be affected at this temperature. air access. This modification could be attributed to The availability of boron oxide could then allow the relatively poor organization of the structure the formation of a glassy phase with various and the initial incorporation of C and o impurities amounts of boron. It is indeed well-known that Because its composition was not changed by the boron significantly decreases the viscosity of glas- ICVI processing of the silicon carbide matrix from ses. The melting temperature of B2O3 is about 480C23 and the eutectic temperature of the pseudo-binary SiO2-B2O3 is, for example, about 450 C24 When considering the flowing ability and link 1 um Fig. 6. Vitreous phase linking the trcated Nicalon m fiber and 10 um the bn interphase (a) deflection at the fiber/BN the BN/matrix and into the Fig. 5. Cracks deflection in a 2D-SiC/BN/SiC composite Fig. 7. Glassy phase in a crack propagating between two tested in tension at 600%C in air neighboring fibers
720 M. Leparoux et al. in boron as compared with nitrogen, that cannot be evidenced by our analyses. 4 Discussion The application of a compliant interphase such as carbon or boron nitride on treated Nicalon@ fibers has been found to improve the mechanical behavior of the Sic/Sic composites at room temperature. The drawbacks of carbon are obvious in oxidizing atmosphere, and BN should exhibit a better oxidation resistance, especially at moderate temperature such as 600°C. However the ICVI-BN deposits always present a turbostratic structure. Even if the temperature of 600°C is low for the onset of hexagonal-BN oxidation, poorer lattice organization might be affected at this temperature. The availability of boron oxide could then allow the formation of a glassy phase with various amounts of boron. It is indeed well-known that boron significantly decreases the viscosity of glasses. The melting temperature of Bz03 is about 480”C23 and the eutectic temperature of the pseudo-binary Si02-B203 is, for example, about 450”C.24 When considering the flowing ability and Fig. 5. Cracks deflection in a 2D-SiC/BN/SiC composite tested in tension at 600°C in air. the volatility of boron oxide, it is not surprising to find such a phase in different parts of the composite, especially in the tows, inside cracks perpendicular or parallel to the fibers (Figs 6 and 7). Sometimes it can even link the cracks walls. The EELS analysis along a decohesion at the fiber/BN interface confirms that the Si, 0 and B elements are present (Section 3.2.2, Analysis of the cracks). The analysis of a crack inside the BN interphase reveals also an increase in oxygen content and the silicon presence with Si-0 bonds formation. From EELS study, it appears that the bulk boron nitride interphase can be modified at 600°C in air. However the decrease of the 0, C and N contents should be examined by further analyses, particularly at different distances from the far end of the cracks, i.e. different conditions of ambient air access. This modification could be attributed to the relatively poor organization of the structure and the initial incorporation of C and 0 impurities. Because its composition was not changed by the ICVI processing of the silicon carbide matrix from Fig. 6. Vitreous phase linking the treated Nicalon@ fiber and the BN interphase. Fig. 7. Glassy phase in a crack propagating between two neighboring fibers
Oxidizing environment infuence on 2D-Sic/ BN/SiC composites hydrogen-methyltrichlorosilane mixture at about access parallel to the fibers, in these very thin sub- 1000C.5 these variations m reorganization resulting fron ight be due to some layers(oxygen diffusion through silica is very slow m an initial interaction it increases with the boron content in glasses but between oxygen from ambient air and C impurity. the amount of boron is low in places that do not Beyond the far end of the cracks, this boron nitride present microcracking) layer could be one way for oxygen access at a low rate inside the composite, especially to the inter faccs with the fibers and the matrix. As shown in the following, the silica formation in place of free carbon should significantly decrease the oxyge Matrix 0um the treated Nicalon fiber and with the sic matrix B-K C-K N-K K ig. 10. EELS spect 22.5 B-K C-K O-K io200 Fig 8. Fracture surface of a 2D-Sic/BN/ SiC tested in ten sion at 600C in air. Rupture occurred at 170 Mpa after Fig. 11. EELS and plasmon (inset) spectra of the treated
Oxidizing environment influence on 2D-SiCIBNjSiC composites 721 a hydrogen-methyltrichlorosilane mixture at about 1000”C,15 these variations might be due to some reorganization resulting from an initial interaction between oxygen from ambient air and C impurity. Beyond the far end of the cracks, this boron nitride layer could be one way for oxygen access at a low rate inside the composite, especially to the interfaces with the fibers and the matrix. As shown in the following, the silica formation in place of free carbon should significantly decrease the oxygen (b) Fig. 8. Fracture surface of a 2D-SiC/BN/SiC tested in tension at 600°C in air. Rupture occurred at 170Mpa after 70 h. access parallel to the fibers, in these very thin sublayers (oxygen diffusion through silica is very slow; it increases with the boron content in glasses but the amount of boron is low in places that do not present microcracking). MatrixJ Fig. 9. Micrograph of the BN interphase and its interface with the treated Nicalon@ fiber and with the SIC matrix. 4, 2ob (,,, 3&I I,,, 4do ,,,, 5do ,,,( 6do ,, Energy Loss (eV) Fig. 10. EELS spectrum of the BN interphase. - Si L2,3 22.5 4 4 5. d- \~ B-K 2 4 8 1 C-K O-K 6 Fig. 11. EELS and plasmon (inset) spectra of the treated Nicalonm fiber/BN interface
M. Leparoux et al Si L2,3 the thinnest Bn interphase could be explained. For the same reasons also when considering the very thin carbon sublayers (when compared to the B-K pyrocarbon interphases)and some possible protec- tion by oxide formers, SiC/BN SiC composites are more oxidation resistant than Sic/C/SiC materials O-K under static loading(Fig. 2 In contrast, the composites with a pyrocarbon interphase exhibit higher resistance to dynamic fatigue in tcnsion-tcnsion, at 600C in air. In that case, the mechanical behavior depends on the slid ing and wear resistance, besides the stability of each constituent in the oxidizing environment and Fig. 12. EELS spectrum of the BN/SiC-matrix interfac the ability for oxygen to diffuse into the composite. Even if it seems that the materials with a bn It is indeed particularly important to examine interphase are less altered by oxidation, it has been accurately the interfaces between the bn inter- shown'2 that the Sic/C/Sic composites present a phase and both the fibers and the matrix. Crack better dynamic fatigue resistance at room tem deflection from mode i to mode il occurs either at per perature. This has been related to the much higher the fiber/BN or BN / matrix interfaces where carbon interfacial shear resistance, twice to five times is the main initial constituent, as shown pre- greater than those of materials with a bn inter viously. 5 The deflection mechanism at the later phase. With regard to the influence of the BN interface was not (or rarely) observed at room thickness, thinner interphases lead to lower sliding temperature owing probably to a weaker fiber/BN resistance and less developed microcracking inside link and different stresses field. When the bn the 0 tows at room temperature. Then, it could be interphase is still bonded with the fibers and the supposed that the fiber/matrix bonding changes matrix, the initial carbon sub-layer has been con- slower during mechanical tests at 600C than those sumed with the formation of silica at the Bn/fiber of thicker interphases. The result is a better reten intcrface. At the BN SiC interface, the initial free tion of the initial mechanical properties.These carbon has been oxidized or transformed in silica hypotheses are based principally on characteristics nd silicon carbide. The csc evaluated at room temperature, espccially the slid con carbide formation clearly leads to strong ing resistance and the crack spacing. A more thor interfacial bondings. When compared to room ough study should consider the coupled influence temperature properties, a worse mechanical beha- of the temperature with the interphase thickness, vior can therefore result the influence of the stress field in the fiber/matrix Obviously, the evolution of the interfacial zones region, and the environmental effects on the inter at a given temperature depends on time but also on facial properties the ability for oxygen to reach the oxidizible layers At this temperature, a sealing of the sic/mati cracks is considered to be negligible, especially far 5 Conclusion from the bn interphase and when a dynamic loading is applied. Thus, the failures of the com- The large potentialities of ceramic composites es appear all the later since the applied stress require the preservation of their thermomech- w as the total crack opening is directly corre- anical properties for long duration, in oxidative to strain. On the other hand, the total environments, under high temperature fat amount of crack opening depends on both cracks conditions. This implies that the fibers, the matrix width and cracks density. A higher density of and the interfacial properties such as load transfer, microcracking might increase the oxidation rate of wear resistance and crack deflection could the interfacial sublayers if we suppose that the maintained gascous diffusion within the radial cracks is not Boron nitride interphases appeared to be more rate-limiting(a reasonable hypothesis at such mod- efficient than pyrocarbon ones under static fatigue erate temperature of 600C). In that casc,morc at 600C in air. Although boron nitride, deposited interfacial zones are damaged when the cracks by ICVI from BCl3-NHy-H2 mixtures, seems to be density increases. Because the crack density decrea- partially oxidized cven at this moderate tempera ses when the BN interphase thickness decreases(at ture, its resistance towards corrosion is still higher least in the range 0.207 um) the better behavior than carbon. As a result of this oxidation, a Si- under static loading of the composites that include b-o glassy phase has been found in the cracks
722 M. Leparoux et al. 7 Si L2,3 J II “‘I’ 1 j n I I 8 b r I1 I I I I I I 1 100 200 300 400 500 600 Energy Loss (eV) Fig. 12. EELS spectrum of the BN/SiC-matrix interface. It is indeed particularly important to examine accurately the interfaces between the BN interphase and both the fibers and the matrix. Crack deflection from mode I to mode II occurs either at the fiber/BN or BN/matrix interfaces where carbon is the main initial constituent, as shown previously.15 The deflection mechanism at the later interface was not (or rarely) observed at room temperature owing probably to a weaker fiber/BN link and different stresses field. When the BN interphase is still bonded with the fibers and the matrix, the initial carbon sub-layer has been consumed with the formation of silica at the BN/fiber interface. At the BNjSiC interface, the initial free carbon has been oxidized or transformed in silica and silicon carbide. These silica or silica plus silicon carbide formation clearly leads to strong interfacial bondings. When compared to room temperature properties, a worse mechanical behavior can therefore result. Obviously, the evolution of the interfacial zones at a given temperature depends on time but also on the ability for oxygen to reach the oxidizible layers. At this temperature, a sealing of the Sic/matrix cracks is considered to be negligible, especially far from the BN interphase and when a dynamic loading is applied. Thus, the failures of the composites appear all the later since the applied stress is low as the total crack opening is directly correlated to strain. On the other hand, the total amount of crack opening depends on both cracks width and cracks density. A higher density of microcracking might increase the oxidation rate of the interfacial sublayers if we suppose that the gaseous diffusion within the radial cracks is not rate-limiting (a reasonable hypothesis at such moderate temperature of 600°C). In that case, more interfacial zones are damaged when the cracks density increases. Because the crack density decreases when the BN interphase thickness decreases (at least in the range 0.24.7pm) the better behavior under static loading of the composites that include the thinnest BN interphase could be explained. For the same reasons, also when considering the very thin carbon sublayers (when compared to the pyrocarbon interphases) and some possible protection by oxide formers, SiC/BN/SiC composites are more oxidation resistant than SiC/C/SiC materials under static loading (Fig. 2). In contrast, the composites with a pyrocarbon interphase exhibit higher resistance to dynamic fatigue in tension-tension, at 600°C in air. In that case, the mechanical behavior depends on the sliding and wear resistance, besides the stability of each constituent in the oxidizing environment and the ability for oxygen to diffuse into the composite. Even if it seems that the materials with a BN interphase are less altered by oxidation, it has been shown’* that the Sic/C/Sic composites present a better dynamic fatigue resistance at room temperature. This has been related to the much higher interfacial shear resistance, twice to five times greater than those of materials with a BN interphase. With regard to the influence of the BN thickness, thinner interphases lead to lower sliding resistance and less developed microcracking inside the 0” tows at room temperature. Then, it could be supposed that the fiber/matrix bonding changes slower during mechanical tests at 600°C than those of thicker interphases. The result is a better retention of the initial mechanical properties.These hypotheses are based principally on characteristics evaluated at room temperature, especially the sliding resistance and the crack spacing. A more thorough study should consider the coupled influence of the temperature with the interphase thickness, the influence of the stress field in the fiber/matrix region, and the environmental effects on the interfacial properties. 5 Conclusion The large potentialities of ceramic composites require the preservation of their thermomechanical properties for long duration, in oxidative environments, under high temperature fatigue conditions. This implies that the fibers, the matrix and the interfacial properties such as load transfer, wear resistance and crack deflection could be maintained. Boron nitride interphases appeared to be more efficient than pyrocarbon ones under static fatigue at 600°C in air. Although boron nitride, deposited by ICVI from BCls-NHs-HZ mixtures, seems to be partially oxidized even at this moderate temperature, its resistance towards corrosion is still higher than carbon. As a result of this oxidation, a SiB-O glassy phase has been found in the cracks
Oxidizing environment infuence on 2D-SiC)BN/SiC composites propagating through the material. The decrease of4.Evans and marshall. d. B. The mechanical beha- the mechanical properties at high temperature is vior of ceramic matrix composites. Acta Metall., 1989, 37 mainly attributed to the evolution of the carbon 256783. 5. Kerans, R.J., Jay,RS. Pagano, N J and Parthasarathy rich interfaces leading to the main formation of T.A. The role of the fiber matrix interface in ceramic silica and consequently to a stronger fiber/matrix composites. Am. Ceram. Soc. Bull, 1989, 68, 429-442 bonding. This evolution can be slowed down by 6. Brennan, J. J, Interfacial studies of chemical-vapor-infil controlling the access of oxygen till the oxidizible trated ceramic matrix composites. Mat. Sci. and Eng, 1990,A126,203-223 ayers. Thinner BN interphase decreases the inter 7. Naslain, R, Fibre-matrix interphases and interfaces in facial shear resistance and mainly the crack den- ceramic matrix composites pro by CVi. Composite sily. Thus, interfacial properties of materials with 8. Lowden. R Schwarz, O.J. and More. K. L small Bn thicknesses, even if they are initially lower are less changed and a better thermomechanical Cerarn. Eng. Sci. Proc, 1993, 14, 375-384 resistance to static and dynamic constraints Bacos, M. P, Carbon-carbon composites: oxidation obtained. The classical Sic/c/SiC materials which sup.J.Phys.I,1993,3,1895-1903 have a much higher interfacial shear resistance still 10. Vandenbulcke, L and Leparoux, M, Silicon and boron exhibit a better behavior under dynamic fatigue at ntaining components by CVD and CVI for high ten- 600C in air, while being more oxidizible. It appears perature ceramic composites. Phys. /V, Colloque C5, 1995,735-751 then of prime importance to increase greatly the ll. Grande, D H, Fibre-matrix bond strength studies of interfacial sliding resistance of SiC/BN/SiC com glass, ceramic, and metal matrix composites. J. Mat. Sci, posites to improve both the static and the dynamic 12. Leparoux, M, Vandenbulcke, L, Goujard, s, Robin- behaviors of such materials. On the other hand it Brosse, C. and Domergue, J M.. Mechanical behavior of would be interesting to prevent the formation of 2D-SiC/BN SiC processed by ICVI. In Proc. Tenth Int carbon-rich interfaces resulting from the fiber Conf on Composite Materials, Vol IV,ed. A. Poursatip and K Street. woodhead, Cambridge, UK, 1995 pp 633-40 treatment and from the sic / matrix infiltration 13. Droillard. C. Elaboration et car Finally, further studies should consider the sites a matrice SiC et a interphase sequence C/SiC. Ph. D influence of the temperature on the interfacial thesis, Bordeaux University, franc propertics and the damaging modes of the materi- 14. Kmetz. M. A. Laliberte. J. M. willis, W. s and Suib S. L, Synthesis, characterization, and tensile strength of als. If the oxygen access parallel to the fibers CVI SiC/BN/ SiC composites. Ceram. Eng. Sci. Pro through the BN interphase itself is important, the 991,12.2161-2174 bn purity, its organization and oxidation resis 15. Leparoux, M, Serin, V, Vandenbulcke, L and Sevely, J TEM and EELS studies of the interphase and interfaces tance should be improved, and multi-sequenced microstructure and chemistry of SiC/BN/SiC composites interphases that include thin BN sub-layers should processed by ICVI. J. Mat Sci., 1997, 32, 4591-4602. CVD and ICVI processes on the boron nitride deposition kinetics. In Proc. of the Xlll Int. Conf. on Chemical Vape Acknowledgements Robinson and R. K. Ulrich, The Electrochem. Society. Pennington, N, 1996, pp 594-599 17. Dambrine, B and Gauthier, G, Role de I'oxydation sur le This research has been supported by the Societe comportement en fatigue du osite Sic/SiC. Revue Europeenne de Propulsion and by the Region des composites et des materiaux avances, 1993, 3, 207-223 Centre through a grant given to M. Leparoux. The 18. Cutard, T, Caracterisation ultrasonore a haute temperature et sous contraintes de traction de composites ceramique authors acknowledge J. Gouineau and G. Lama zouade from SEP for the mechanical characteriza- /9 ceramique Ph. D thesis, Limoges University, France, 1993 tions and for the SEM observations, respectively des fibres et de la zone interfacial fibre matrice sur le co portement mecanique des composites Sic/C/SiC et SiC/ AS-L apres vieillissement thermique sous atmos controlee, Ph D. thesis, Bordeaux University, France, 199 References 20. Stephan, O, Ajayan, P. M, Colliex, C,, Redlich, P. H Lambert, J M, Bernier, P and Lefin, P,, Doping graphite I. Jouin, J. M., Cotteret, J and Christin, F. Sic-Sic inter- and carbon nanotube structures with boron and nitrogen hases case history. In Proceeding s of the 2nd europe Science,1994,266,l683-1685 Designing Ceramic Interfaces: Under- 21. Egerton, R. F,, Electron Energy-Loss Spectroscopy in the anding and tailoring interfaces for Coating, Composite 22. Kun ron Microscope. Plenum Press, New York, 1986 and Joining Applications, ed S. D. Peteves, CEE, Petten, 22. Kundman, M, EL/ P Software Instruction Manual. Gatan Netherlands, 1 1-13 November 1991, pp. 191-204 c,1993 CVD/CVI Pulse aux materiaux composites thermostructuraux. Ph. D thesis nology, Vol 3, Viscosity and Relaxation, ec Bordeaux University, france, 1993 nann and n.. Kreidl. Academic press. Inc. Orlando 3. Leparoux, M, Vandenbulcke, L. and Clinard, C, Influ 1986,p.245 ence of isothermal CVD and Cvi processes on the 24. Rockett, T. J. and Foster. W. R. Phase relations in the nitride deposition kinctics and microstructure, Submitted system boron oxide silica. J. Am. Ceram. Soc, 1965, 48, to j. Am. Ceram. sc 7580
Oxidizing environment influence on ZD-SiCIBNjSiC composites 123 propagating through the material. The decrease of the mechanical properties at high temperature is mainly attributed to the evolution of the carbonrich interfaces leading to the main formation of silica and consequently to a stronger fiber/matrix bonding. This evolution can be slowed down by controlling the access of oxygen till the oxidizible layers. Thinner BN interphase decreases the interfacial shear resistance and mainly the crack density. Thus, interfacial properties of materials with small BN thicknesses, even if they are initially lower, are less changed and a better thermomechanical resistance to static and dynamic constraints is obtained. The classical Sic/C/Sic materials which have a much higher interfacial shear resistance still exhibit a better behavior under dynamic fatigue at 600°C in air, while being more oxidizible. It appears then of prime importance to increase greatly the interfacial sliding resistance of SiC/BN/SiC composites to improve both the static and the dynamic behaviors of such materials. On the other hand, it would be interesting to prevent the formation of carbon-rich interfaces resulting from the fiber treatment and from the Sic/matrix infiltration. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. Finally, further studies should consider the influence of the temperature on the interfacial properties and the damaging modes of the materials. If the oxygen access parallel to the fibers through the BN interphase itself is important, the BN purity, its organization and oxidation resistance should be improved, and multi-sequenced interphases that include thin BN sub-layers should be studied. 14. 15. 16. Acknowledgements 17. This research has been supported by the Societe Europeenne de Propulsion and by the Region Centre through a grant given to M. Leparoux. The authors acknowledge J. Gouineau and G. Lamazouade from SEP for the mechanical characterizations and for the SEM observations, respectively. 18. 19 References 20. Jouin, J. M., Cotteret, J. and Christin, F., SiCSiC interphases case history. In Proceedings of the 2nd European Colloquium on Designing Ceramic Interfaces: Understanding and tailoring interfaces for Coating, Composite and Joining Applications, ed. S. D. Peteves, CEE, Petten, Netherlands, 11-13 November 1991, pp. 191-204. Dupel, P., CVDjCVI pulsee du pyrocarbone: application aux materiaux composites thermostructuraux. Ph.D. thesis, Bordeaux University, France, 1993. Leparoux, M., Vandenbulcke, L. and Clinard, C., Influence of isothermal CVD and CVI processes on the boron nitride deposition kinetics and microstructure. Submitted to J. Am. Ceram. Sot. 21. 22. 23. 24. Evans, A. G. and Marshall, D. B., The mechanical behavior of ceramic matrix composites. Acta Metall., 1989, 37, 256783. Kerans, R. J., Jay, R. S., Pagano, N. J. and Parthasarathy, T. A., The role of the fiber matrix interface in ceramic composites. Am. Ceram. Sot. Bull, 1989, 68, 429-442. Brennan, J. J., Interfacial studies of chemical-vapor-infiltrated ceramic matrix composites. Mat. Sci. and Eng, 1990, A126,203-223. Naslain, R., Fibre-matrix interphases and interfaces in ceramic matrix composites processed by CVI. Composite Interfaces, 1993, 1, 253-286. Lowden, R. A., Schwarz, 0. J. and More, K. L., Improved fiber coatings for Nicalon/SiC composites. Ceram. Eng. Sci. Proc, 1993, 14, 375-384. Bacos, M. P., Carbon-carbon composites: oxidation behavior and coatings protection. J. Phys. IV, ~011. C7, supl. J. Phys. ZZI, 1993, 3, 1895-1903. Vandenbulcke, L. and Leparoux, M., Silicon and boron containing components by CVD and CVI for high temperature ceramic composites. J. Phys. IV, Colloque C5, 1995, 735-751. Grande, D. H., Fibre-matrix bond strength studies of glass, ceramic, and metal matrix composites. J. Mat. Sci, 1988,23, 311-328. Leparoux, M., Vandenbulcke, L., Goujard, S., RobinBrosse, C. and Domergue, J. M., Mechanical behavior of 2D-SiC/BN/SiC processed by ICVI. In Proc. Tenth Znt. Conf on Composite Materials, Vol. IV, ed. A. Poursatip and K. Street. Woodhead, Cambridge, UK, 1995 pp. 63340. Droillard, C., Elaboration et caracterisation de composites d matrice Sic et a interphase sequencee C/Sic. Ph.D. thesis, Bordeaux University, France, 1993. Kmetz, M. A., Laliberte, J. M., Willis, W. S. and Suib, S. L., Synthesis, characterization, and tensile strength of CVI SiC/BN/SiC composites. Ceram. Eng. Sci. Proc, 1991, 12, 2161-2174. Leparoux, M., Serin, V., Vandenbulcke, L. and Sevely, J., TEM and EELS studies of the interphase and interfaces microstructure and chemistry of SiC/BN/SiC composites processed by ICVI. J. Mat. Sci., 1997, 32, 45914602. Leparoux, M. and Vandenbulcke, L., Influence of the CVD and ICVI processes on the boron nitride deposition kinetics. In Proc. of the XIII Int. Conf. on Chemical Vapor Deposition, ed. T. M. Besmann, M. D. Allendorf, McD. Robinson and R. K. Ulrich, The Electrochem. Society, Pennington, NJ, 1996, pp. 594599. Dambrine, B. and Gauthier, G., Role de l’oxydation sur le comportement en fatigue d’un composite Sic/Sic. Revue des composites et des materiaux avancks, 1993, 3, 2077223. Cutard, T., Caracterisation ultrasonore a haute temperature et sous contraintes de traction de composites ceramique/ ceramique. Ph.D. thesis, Limoges University, France, 1993. Labrugere, C., Influence de l’evolution physico-chimique des fibres et de la zone interfaciale fibre/matrice sur le comportement mecanique des composites Sic/C/Sic et Sic/ MAS-L apres vieillissement thermique sous atmosphere controlee. Ph.D. thesis, Bordeaux University, France, 1994. Stephan, O., Ajayan, P. M., Colliex, C., Redlich, P. H., Lambert, J. M., Bernier, P. and Lefin, P., Doping graphite and carbon nanotube structures with boron and nitrogen. Science, 1994, 266, 1683-1685. Egerton, R. F., Electron Energy-Loss Spectroscopy in the Electron Microscope. Plenum Press, New York, 1986. Kundman, M., EL/P Software Instruction Manual. Gatan Inc., 1993. Scholze, H. and Kreidl, N. J., Glass-Science and Technology, Vol. 3, Viscosity and Relaxation, ed. D. R. Uhlmann and N. J. Kreidl, Academic Press, Inc., Orlando, FL, 1986, p. 245. Rockett, T. J. and Foster, W. R., Phase relations in the system boron oxide-silica. J. Am. Ceram. Sot, 1965, 48, 7580
