Availableonlineatwww.sciencedirect.co Science direct materials letters ELSEVIER Materials Letters 60(2006)3197-320 www.elseviercom/locate/matlet Oxidation behavior of 3D Hi-Nicalon/SiC composite Shoujun Wu", Aifei Cheng, Litong Zhang, Yongdong Xu, Qing Zhang National Key Laboratory of Thermostructure Composite Materials, Northwestern Pol al University, Xi'an Shaanxi 710072, People's Republic of China Available online 20 March 2006 Abstract Oxidation behavior of a three dimensional (3D) Hi-Nicalon/SiC composite with Cvd Sic coating was investigated in the simulated air using a thermogravimetric analysis(TGA)device. Below 1100C, the oxidation kinetics was controlled by gas diffusion through the defects in the Sic matrix and coating and resulted in the consumption of PyC interphase. The residual flexural strength did have not a remarkable fluctuation and the relationship between the residual strength to temperature and weight change to temperature of the 3D Hi-Nicalon/PyC/SiC composite indicated the same regularity. Above 1200C, the oxidation kinetics was controlled by oxygen diffusion through the SiO2 scale formed on the Sic coating and matrix. And the residual flexural strength of the composites was governed by the strength degradation of the Hi-Nicalon fiber. After oxidation, the fracture displacement in flexural tests increased with the weight loss increasing and the fracture mode showed a non-brittle pattern C 2006 Elsevier B V. All rights reserved. Keywords: Hi-Nicalon/PyC/SiC; Oxidation; Coating: Residual flexural strength 1. Introduction and the strength degradation of the fiber [8]. Thus, for Nicalon/ SiC composites, the degradation of mechanical properties after Silicon carbide fiber reinforced silicon carbide composites short periods of exposure to air was due to the oxidation and (SiC/SiC)exhibit better oxidation resistance than C/SiC, and are removal of the PyC interlayer [9]. Hi-Nicalon is oxygen-free considered as one of the most promising structural materials for fibers consisting of a mixture of Sic-nanocrystals( 5 nm in high temperature applications [1-3]. In order to improve the mean size)and free carbon [C/Si(at)ratio=1.39]. They do not mechanical properties of fiber reinforced silicon carbide matrix undergo decomposition at high temperatures since they do not composites, interphase of compliant material with low shear contain significant amount of SiCoy phase [10, 11].Hi- strength is necessary. It has also been recognized that the Nicalon/SiC composites are expected to be used above 1400C, interphase layers should be deposited parallel to the fiber and are being paid more and more attention [12]. However the surface and weakly bonded to one another, and the interphase oxidation behavior, especially the oxidation-induced mechan- should be strongly bonded to the fiber surface [2, 4]. Pyrocarbon ical properties changes of C-interphase Hi-Nicalon/SiC(Hi- (PyC)and hexagonal-BN(hex-BN) have been widely used as Nicalon/PyC/SiC) during oxidation have not been much the interphase materials [2, 4-6]. PyC is much liable to researched oxidation than hex-BN. However, hex-BN was poor bonded In order to deeply understanding the main oxidation to the Sic fiber surface and not sufficiently better to be viable at mechanisms of Hi-Nicalon/Py C/Sic materials and the effects above1300°C[4,7 of PyC interphase on the flexural strength after oxidation in air, The Nicalon fiber is consist of Sic-nanocrystals(1-2 nm in the oxidation behavior of a 3D Hi-Nicalon/Py C/SiC with a size)and free carbon embedded in an amorphous SiCO CVD Sic coating was investigated in the present paper. The matrix. At the oxidizing atmosphere, oxidation and decompo- weight change of the composite during oxidation was sition of SicxOy result in pores produced on the fiber surface monitored with thermogravimetric analysis (TGA)device And the relationship between the residual flexural strength to Corresponding author. Tel: +8629 8848 6068 828; fax: +8629 8849 4620. temperature and the weight change to temperature of the E-mailaddress:shoujun-wu(@163.com(S.Wu) composite were also researched. 0167-577X/S-see front matter o 2006 Elsevier B V. All rights reserved. doi:10.1016/ malet200602.072
Oxidation behavior of 3D Hi–Nicalon/SiC composite Shoujun Wu ⁎, Laifei Cheng, Litong Zhang, Yongdong Xu, Qing Zhang National Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, Xi'an Shaanxi 710072, People's Republic of China Received 18 September 2005; accepted 22 February 2006 Available online 20 March 2006 Abstract Oxidation behavior of a three dimensional (3D) Hi–Nicalon/SiC composite with CVD SiC coating was investigated in the simulated air using a thermogravimetric analysis (TGA) device. Below 1100 °C, the oxidation kinetics was controlled by gas diffusion through the defects in the SiC matrix and coating and resulted in the consumption of PyC interphase. The residual flexural strength did have not a remarkable fluctuation and the relationship between the residual strength to temperature and weight change to temperature of the 3D Hi–Nicalon/PyC/SiC composite indicated the same regularity. Above 1200 °C, the oxidation kinetics was controlled by oxygen diffusion through the SiO2 scale formed on the SiC coating and matrix. And the residual flexural strength of the composites was governed by the strength degradation of the Hi–Nicalon fiber. After oxidation, the fracture displacement in flexural tests increased with the weight loss increasing and the fracture mode showed a non-brittle pattern. © 2006 Elsevier B.V. All rights reserved. Keywords: Hi–Nicalon/PyC/SiC; Oxidation; Coating; Residual flexural strength 1. Introduction Silicon carbide fiber reinforced silicon carbide composites (SiC/SiC) exhibit better oxidation resistance than C/SiC, and are considered as one of the most promising structural materials for high temperature applications [1–3]. In order to improve the mechanical properties of fiber reinforced silicon carbide matrix composites, interphase of compliant material with low shear strength is necessary. It has also been recognized that the interphase layers should be deposited parallel to the fiber surface and weakly bonded to one another, and the interphase should be strongly bonded to the fiber surface [2,4]. Pyrocarbon (PyC) and hexagonal–BN (hex–BN) have been widely used as the interphase materials [2,4–6]. PyC is much liable to oxidation than hex–BN. However, hex–BN was poor bonded to the SiC fiber surface and not sufficiently better to be viable at above 1300 °C [4,7]. The Nicalon fiber is consist of SiC–nanocrystals (1∼2 nm in size) and free carbon embedded in an amorphous SiCxOy matrix. At the oxidizing atmosphere, oxidation and decomposition of SiCxOy result in pores produced on the fiber surface and the strength degradation of the fiber [8]. Thus, for Nicalon/ SiC composites, the degradation of mechanical properties after short periods of exposure to air was due to the oxidation and removal of the PyC interlayer [9]. Hi–Nicalon is oxygen–free fibers consisting of a mixture of SiC–nanocrystals (≈5 nm in mean size) and free carbon [C/Si (at) ratio = 1.39]. They do not undergo decomposition at high temperatures since they do not contain significant amount of SiCxOy phase [10,11]. Hi– Nicalon/SiC composites are expected to be used above 1400 °C, and are being paid more and more attention [12]. However the oxidation behavior, especially the oxidation–induced mechanical properties changes of C–interphase Hi–Nicalon/SiC (Hi– Nicalon/PyC/SiC) during oxidation have not been much researched. In order to deeply understanding the main oxidation mechanisms of Hi–Nicalon/PyC/SiC materials and the effects of PyC interphase on the flexural strength after oxidation in air, the oxidation behavior of a 3D Hi–Nicalon/PyC/SiC with a CVD SiC coating was investigated in the present paper. The weight change of the composite during oxidation was monitored with thermogravimetric analysis (TGA) device. And the relationship between the residual flexural strength to temperature and the weight change to temperature of the composite were also researched. Materials Letters 60 (2006) 3197–3201 www.elsevier.com/locate/matlet ⁎ Corresponding author. Tel.: +86 29 8848 6068 828; fax: +86 29 8849 4620. E-mail address: shoujun_wu@163.com (S. Wu). 0167-577X/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2006.02.072
3198 S. Wu et al. / Materials Letters 60(2006)3197-3201 2. Experimental procedure Interphase layer 2.1. Preparation of 3D Hi-Nicalon/PyC/SiC composite specimens Hi-NicalonTM silicon carbide fiber from Japan Nippon Matrix micro crack Carbon was employed. The fiber preform was prepared using four-step three dimensional (4-step 3D) braiding method, and was supplied by the Nanjing Institute of Glass Fiber, Peoples epublic of China. Low pressure chemical vapor infiltration (LPCVD) process was employed to deposit pyrolytic carbon (PyC) interphase and the silicon carbide matrix. The volume Fig. 2. Fine matrix micro cracks in the 3D Hi-Nicalon/PyC/SiC composite. fraction of fibers was about 40% and the braiding angle was about 20. The interfacial layer of PyC was deposited for I h at 870C and 5 kPa with C3H6. The deposited PyC interphase mal Analysis System(thermobalance with a sensitivity of layer is about 0.2 um Methyltrichlorosilane (MTS, CH3SiCl3) +0.0001 mg) for 900 min. The simulated air was a mixture of was used for the deposition of the SiC matrix. MTS vapor was oxygen and argon with a volume ratio of 22 to 78. The gas flow carried by bubbling hydrogen. The conditions for deposition of speed was about 3. 5 cms. Weight changes related to oxidation SiC matrix were as follows: the deposition temperature was time of specimens was recorded with TG mode. The oxidation 1100C, pressure was 5 kPa, time was 120 h, the molar ratio of test procedure was as follows: Firstly, the specimen was hold for H2 to methyltrchlorosilane (MTS)was 10. Argon was employed 15 min at 25C in argon. Secondly, heat up the specimen in as the dilute gas to slow down the chemical reaction rate of argon with controlled heating rate. Finally, introduce the deposition. The dimension of as-received SiC/SiC sample was simulated air at the desired temperature and start oxidation 2.5x42x 30.0 mm. The Sic coating was prepared on the process specimens for 20 h to seal the open ends of the fibers after cutting from the prepare 2.3. Measurements of the composite specimens 2.2. Oxidation tests The flexural strength of the specimens before and after oxidation was measured by a three-point bending method, The oxidation of the Hi-Nicalon/PyC/SiC in a simulated which was carried out on an Instron 1195 machine at room air was conducted in METTLER TOLEDO STAR Ther- temperature. The span dimension was 20 mm and the loading rate was 0.5 mm/min. The fracture sections and the surfaces of the specimens were observed on a scanning electron microscope (SEM,S4700). 3. Results and discussion 3.1. Microstructure of the 3D Hi-Nicalon/Py C/SiC composite Fig. I showed the surface and interior section morphologies of the 3D Hi-Nicalon/PyC/SiC composite Cracks were neither observed in the outer CVD SiC coating, nor in the SiC matrix. However, finely micro cracks could be observed in the matrix for the Hi-Nicalon/PyC/ Unsealed pore Fig. 1. Surface and interior section morphologies of the 3D Hi-Nicalon/PyC Fig 3. Unsealed pores in CVD SiC coating
2. Experimental procedure 2.1. Preparation of 3D Hi–Nicalon/PyC/SiC composite specimens Hi–Nicalon™ silicon carbide fiber from Japan Nippon Carbon was employed. The fiber preform was prepared using four–step three dimensional (4–step 3D) braiding method, and was supplied by the Nanjing Institute of Glass Fiber, People’s Republic of China. Low pressure chemical vapor infiltration (LPCVI) process was employed to deposit pyrolytic carbon (PyC) interphase and the silicon carbide matrix. The volume fraction of fibers was about 40% and the braiding angle was about 20°. The interfacial layer of PyC was deposited for 1 h at 870 °C and 5 kPa with C3H6. The deposited PyC interphase layer is about 0.2 μm. Methyltrichlorosilane (MTS, CH3SiCl3) was used for the deposition of the SiC matrix. MTS vapor was carried by bubbling hydrogen. The conditions for deposition of SiC matrix were as follows: the deposition temperature was 1100 °C, pressure was 5 kPa, time was 120 h, the molar ratio of H2 to methyltrchlorosilane (MTS) was 10. Argon was employed as the dilute gas to slow down the chemical reaction rate of deposition. The dimension of as–received SiC/SiC sample was 2.5 × 4.2 × 30.0 mm. The SiC coating was prepared on the specimens for 20 h to seal the open ends of the fibers after cutting from the prepared composite. 2.2. Oxidation tests The oxidation of the Hi-Nicalon/PyC/SiC in a simulated air was conducted in METTLER TOLEDO STARe Thermal Analysis System (thermobalance with a sensitivity of ± 0.0001 mg) for 900 min. The simulated air was a mixture of oxygen and argon with a volume ratio of 22 to 78. The gas flow speed was about 3.5 cm·s–1 . Weight changes related to oxidation time of specimens was recorded with TG mode. The oxidation test procedure was as follows: Firstly, the specimen was hold for 15 min at 25 °C in argon. Secondly, heat up the specimen in argon with controlled heating rate. Finally, introduce the simulated air at the desired temperature and start oxidation process. 2.3. Measurements of the composite specimens The flexural strength of the specimens before and after oxidation was measured by a three–point bending method, which was carried out on an Instron 1195 machine at room temperature. The span dimension was 20 mm and the loading rate was 0.5 mm/min. The fracture sections and the surfaces of the specimens were observed on a scanning electron microscope (SEM, S4700). 3. Results and discussion 3.1. Microstructure of the 3D Hi–Nicalon/PyC/SiC composite Fig. 1 showed the surface and interior section morphologies of the 3D Hi–Nicalon/PyC/SiC composite. Cracks were neither observed in the outer CVD SiC coating, nor in the SiC matrix. However, finely micro cracks could be observed in the matrix for the Hi–Nicalon/PyC/ Fig. 1. Surface and interior section morphologies of the 3D Hi–Nicalon/PyC/ SiC composite (a) Surface; (b) interior section. Fig. 2. Fine matrix micro cracks in the 3D Hi–Nicalon/PyC/SiC composite. Fig. 3. Unsealed pores in CVD SiC coating. 3198 S. Wu et al. / Materials Letters 60 (2006) 3197–3201
S. Wu et al. Materials Letters 60(2006)3197-3201 3199 1.8 Cvd SiC 200400600800100012001400 Fig. 4. Relations of CTE to temperatures for the 3D Hi-Nicalon/PyC/SIC composite and CVD Sic SiC composite under higher magnification as shown in Fig. 2. Between the fiber and matrix a Pyc interphase layer could be observed and gaps were scarcely observed either between the fiber and the interphase or between the matrix and the interphase. Moreover, some pores existed the Cvd sic coating as shown in Fig 3. As for fiber reinforced ceramic matrix composites, cracks were resulted from the mismatch of thermal expansion coefficient(CTE between the composite and coating, and between the fibers and matrix, respectively. The CtE of Hi-Nicalon SiC fiber and SiC matrix w about3.1~3.5×10-6Cand4.6×10-°C, respectively[2,10,13]Fig. 4 showed the relations of CTE to temperatures of the 3D Hi-Nicalon/ PyC/Sic composite and CVD SiC from 25 C to 1400oC. The relations Fig. 6. Morphology of nonuniform consumption of PyC phase in Hi-Nicalon/ of CTE to temperature of Hi-Nicalon/Py C/SiC composite have a PyC/SiC composite at inat:(a)600°C;(b) similar tendency to those of CVD SiC and the difference between them was very little. Thus, a little tensile stress was produced in the Sic coating or matrix during cooling process and the thermal cracks were scarcely produced in them. Some matrix pores formed between the fiber bundles were open on the surface of the composite after machined. Thus, open pores could be observed in the SiC coating 3.2. Oxidation behavior of the 3D Hi-Nicalon/Py C/SiC composite Fig. 5 showed the TGA results of the Hi-Nicalon/PyC/SiC composite after oxidation in the simulated air for 900 min. Below 55065075085095010501150125013501450 1100C, oxidation of the composite led to a weight loss. Above 1100C, oxidation resulted in a weight gain in a parabolic-linear 600 Time(min) 5006007008009001000I1001200130014001500 Fig. 5.(a)Relations of weight changes to temperature of the 3D Hi-Nicalon/ PyC/SiC composite after oxidized for 900 min(b) Typical TG curves during the Fig. 7. Relations of residual strength to tures of the Hi-Nicalon/PyC/SiC first 400 min composite after oxidized for 900 min
SiC composite under higher magnification as shown in Fig. 2. Between the fiber and matrix a PyC interphase layer could be observed and gaps were scarcely observed either between the fiber and the interphase or between the matrix and the interphase. Moreover, some pores existed in the CVD SiC coating as shown in Fig. 3. As for fiber reinforced ceramic matrix composites, cracks were resulted from the mismatch of thermal expansion coefficient (CTE) between the composite and coating, and between the fibers and matrix, respectively. The CTE of Hi–Nicalon SiC fiber and SiC matrix was about 3.1∼3.5 × 10–6 /°C and 4.6 × 10–6 /°C, respectively [2,10,13]. Fig. 4 showed the relations of CTE to temperatures of the 3D Hi–Nicalon/ PyC/SiC composite and CVD SiC from 25 °C to 1400 °C. The relations of CTE to temperature of Hi–Nicalon/PyC/SiC composite have a similar tendency to those of CVD SiC and the difference between them was very little. Thus, a little tensile stress was produced in the SiC coating or matrix during cooling process and the thermal cracks were scarcely produced in them. Some matrix pores formed between the fiber bundles were open on the surface of the composite after machined. Thus, open pores could be observed in the SiC coating. 3.2. Oxidation behavior of the 3D Hi–Nicalon/PyC/SiC composite Fig. 5 showed the TGA results of the Hi–Nicalon/PyC/SiC composite after oxidation in the simulated air for 900 min. Below 1100 °C, oxidation of the composite led to a weight loss. Above 1100 °C, oxidation resulted in a weight gain in a parabolic–linear manner. Fig. 4. Relations of CTE to temperatures for the 3D Hi–Nicalon/PyC/SiC composite and CVD SiC. Fig. 5. (a) Relations of weight changes to temperature of the 3D Hi–Nicalon/ PyC/SiC composite after oxidized for 900 min (b) Typical TG curves during the first 400 min. Fig. 6. Morphology of nonuniform consumption of PyC phase in Hi–Nicalon/ PyC/SiC composite after oxidized for 900 min at : (a) 600 °C; (b) 900 °C. Fig. 7. Relations of residual strength to temperatures of the Hi–Nicalon/PyC/SiC composite after oxidized for 900 min. S. Wu et al. / Materials Letters 60 (2006) 3197–3201 3199
S. Wu et al. / Materials Letters 60(2006)3197-3201 Below the deposition temperature of SiC, defects in the coating and matrix as shown in Figs. 2 and 3, played a role of transporting channels for oxygen diffusion inside. It was clearly that these channels were too narrow for oxygen to diffuse freely and oxidation only took place along the PyC interphase. The oxidation kinetics was controlled by gas diffusion through these channels consisting of defects in the coating and matrix, and resulting in the consumption of Pyc phase as shown i Fig 6 Below 800C, the oxidation weight loss nearly kept constant. From 800 oC to 1000 oC. the width of the microcracks in Sic matrix decreased as temperature increased, but the oxidation reaction rate obeyed Arrehnius equation. Moreover, the oxidation of Sic was very slow [14]. At the specific oxidation condition (i.e. temperature, oxidizing atmosphere, and reactant), the oxidation rate is a constant. for 900 min Fig 9. Silica formed between the SiC fibers and matrix after oxidized at 1400C Consequently, the weight loss increased with temperatu a linear form to oxidation time. Above 1000 C. though the defects changed after oxidation. The failure behavior of the as prepared loss, oxidation of Sic became more and more significant as composite was rather brie and exhibited seep stress drops atter the temperature increased. Thus, the weight loss decreased with temper- very gradual after the maximum stress point Above 1100C, the width of micro cracks became very narrow due It is well known that the strength of fiber reinforced ceramic matrix composites is greatly influenced by the strength of the reinforcing to the thermal expansion, and the oxidation of Sic became significant fibers, the characteristics of fiber/matrix interface and the residual stresses caused by thermal expansion mismatch between fibers and diffusion channels. At the beginning, oxidation of Sic resulted in the matrix [2, 15, 16]. Usually a proper weak interphase is favorable. The E of Hi-Nicalon SiC fiber and Sic matrix was about 31-3.5x formation of protective SiO2 and rapid weight gain in a parabolic 10-6/ oC and 4.6x10-6/C, respectively [2, 10,13].Hence, the manner. With the oxidation processed, oxygen diffusion was slow compressive stress within the interfacial phase along the fiber radial down and balanced intensively. Consequently, the weight gain rate kept direction was generated after the composite was cooled down from the nearly constant as the oxidation proceeded further. 3.3. Relations behveen residual flexural strength and weight change for the Hi-Nicalon SiC fiber to debond and to be pulled out from the silicon carbide matrix. Because oxidations took place along and weaken the interphase, the Hi-Nicalon SiC fiber could be much easier Fig. 7 showed the relationship between residual strength and to debond. The more the interphase consumed, the easier the fiber temperatures of the 3D Hi-Nicalon/Pyc/SiC composite oxidized in the debond. The consumption of carbon interphase resulted in a weak simulated air at different temperatures for 900 min. Below 1200C, the interphase bond between fiber and matrix, and then led to the residual Sationshin between residual strength and weight change of the 3D flexural strength wi油ham小 ange as the weight loss indicated the same regularity. The residual flexural strength of the degradation with temperature increased [1o, I. Furthermore, oxygen opposites increased with weight loss decreasing in oxidation. At 1000oC, the weight loss arrived at maximum value, and the residual between fibres and matrix became very small after the PyC interphase flexural strength got the minimum value accordingly. Above 1200C, has been consumed by oxidation. Thus, on the one hand, the e residual flexural strength of the 3D Hi-Nicalon/PyC/SiC composite interphase bond between fiber and matrix was weakened; on the other decreased with oxidation temperature increase. Several load-displacement curves of the composite were shown in hand, it caused fiber degradation on mechanical strength [8 As a result, the residual flexural strength of the composite st Fig.8. It can be seen that the failure behavior of the composite was decreased as temperature increased. And the failure behavior for the oxidized composite changed from brittle to a non-brittle pattern. 600°C 4. Conclusions 1. The oxidation behavior of the 3D Hi-Nicalon/PyC/SiC composite has been investigated in simulated air, on the basis of weight and flexural strength change. Below 1100C 300 the oxidation kinetics was controlled by gas diffusion through defects in the Sic matrix and coating. Above 1200C, the oxidation kinetics was controlled by oxygen diffusion through Displacement( mm) 2. Below 1200C, the residual flexural strength did not have a Fig 8. Several load-displat curves of the Hi-Nicalon/PyC/SiC composite remarkable fluctuation, and the relationship between residual strength and weight change of the 3D Hi-Nicalon/PyC/Sic
Below the deposition temperature of SiC, defects in the coating and matrix as shown in Figs. 2 and 3, played a role of transporting channels for oxygen diffusion inside. It was clearly that these channels were too narrow for oxygen to diffuse freely and oxidation only took place along the PyC interphase. The oxidation kinetics was controlled by gas diffusion through these channels consisting of defects in the coating and matrix, and resulting in the consumption of PyC phase as shown in Fig. 6. Below 800 °C, the oxidation weight loss nearly kept constant. From 800 °C to 1000 °C, the width of the microcracks in SiC matrix decreased as temperature increased, but the oxidation reaction rate obeyed Arrehnius equation. Moreover, the oxidation of SiC was very slow [14]. At the specific oxidation condition (i.e. temperature, oxidizing atmosphere, and reactant), the oxidation rate is a constant. Consequently, the weight loss increased with temperature and showed a linear form to oxidation time. Above 1000 °C, though the defects were still open and oxygen could diffuse inward and resulted in weight loss, oxidation of SiC became more and more significant as temperature increased. Thus, the weight loss decreased with temperature increased. Above 1100 °C, the width of micro cracks became very narrow due to the thermal expansion, and the oxidation of SiC became significant and was controlled by oxygen diffusion through SiO2 scale. Furthermore, the volume expansion of SiO2 would enclose the oxygen diffusion channels. At the beginning, oxidation of SiC resulted in the formation of protective SiO2 and rapid weight gain in a parabolic manner. With the oxidation processed, oxygen diffusion was slow down and balanced intensively. Consequently, the weight gain rate kept nearly constant as the oxidation proceeded further. 3.3. Relations between residual flexural strength and weight change Fig. 7 showed the relationship between residual strength and temperatures of the 3D Hi–Nicalon/PyC/SiC composite oxidized in the simulated air at different temperatures for 900 min. Below 1200 °C, the residual flexural strength did not have a remarkable fluctuation and the relationship between residual strength and weight change of the 3D Hi–Nicalon/PyC/SiC composite oxidized at different temperature indicated the same regularity. The residual flexural strength of the composites increased with weight loss decreasing in oxidation. At 1000 °C, the weight loss arrived at maximum value, and the residual flexural strength got the minimum value accordingly. Above 1200 °C, the residual flexural strength of the 3D Hi–Nicalon/PyC/SiC composite decreased with oxidation temperature increase. Several load–displacement curves of the composite were shown in Fig. 8. It can be seen that the failure behavior of the composite was changed after oxidation. The failure behavior of the as prepared composite was rather brittle and exhibited steep stress drops after the maximum stress point. After oxidation for 900 min, the stress drop was very gradual after the maximum stress point. It is well known that the strength of fiber reinforced ceramic matrix composites is greatly influenced by the strength of the reinforcing fibers, the characteristics of fiber/matrix interface and the residual stresses caused by thermal expansion mismatch between fibers and matrix [2,15,16]. Usually a proper weak interphase is favorable. The CTE of Hi–Nicalon SiC fiber and SiC matrix was about 3.1∼3.5 × 10– 6 /°C and 4.6 × 10– 6 /°C, respectively [2,10,13]. Hence, the compressive stress within the interfacial phase along the fiber radial direction was generated after the composite was cooled down from the infiltration temperature (1100 °C) to room temperature. It was difficult for the Hi–Nicalon SiC fiber to debond and to be pulled out from the silicon carbide matrix. Because oxidations took place along and weaken the interphase, the Hi–Nicalon SiC fiber could be much easier to debond. The more the interphase consumed, the easier the fiber debond. The consumption of carbon interphase resulted in a weak interphase bond between fiber and matrix, and then led to the residual flexural strength with a slight fluctuation change as the weight loss. Above 1200 °C, the strength of Hi–Nicalon fibers exhibited degradation with temperature increased [10,11]. Furthermore, oxygen could diffuse into the composite and resulted in silica locally formed between the fibers and matrix as shown in Fig. 9. The adhesion between fibres and matrix became very small after the PyC interphase has been consumed by oxidation. Thus, on the one hand, the interphase bond between fiber and matrix was weakened; on the other hand, it caused serious fiber degradation on mechanical strength [8]. As a result, the residual flexural strength of the composite strongly decreased as temperature increased. And the failure behavior for the oxidized composite changed from brittle to a non–brittle pattern. 4. Conclusions 1. The oxidation behavior of the 3D Hi–Nicalon/PyC/SiC composite has been investigated in simulated air, on the basis of weight and flexural strength change. Below 1100 °C, the oxidation kinetics was controlled by gas diffusion through defects in the SiC matrix and coating. Above 1200 °C, the oxidation kinetics was controlled by oxygen diffusion through the SiO2 scale formed on the SiC coating and matrix. 2. Below 1200 °C, the residual flexural strength did not have a remarkable fluctuation, and the relationship between residual strength and weight change of the 3D Hi–Nicalon/PyC/SiC Fig. 8. Several load–displacement curves of the Hi–Nicalon/PyC/SiC composite in flexural strength tests. Fig. 9. Silica formed between the SiC fibers and matrix after oxidized at 1400 °C for 900 min. 3200 S. Wu et al. / Materials Letters 60 (2006) 3197–3201
S. Wu et al. Materials Letters 60(2006)3197-3201 composite oxidized at different temperature indicated the 3]P Fenici, A.J. Frias Rebelo, R H. Jones, A Kohyama, LL. Snead, J NucL. same regularity. The residual flexural strength of the Mate:.258-263(1998)215 composite was governed by the consumption of Pyc [5]S. Jacques, A. Lopez-Marure, C. Vincent, H. Vincent, J. Bouix, I.Eur. interphase, and the oxidation had few effects on the strength of the composite. Above 1200oC, the residual flexural [6]RJ. Kerans, R.S. Hay, T.A. Parthasarathy, Curr Opin. Solid State Mater. strength of the composites was governed by the strength Sc.4(1999)445 degradation of the Hi-Nicalon fib [7RE Tressler, Composites, Part A 30(1999)429. 3. Because oxidation took place along and weakened the [8]M. Takeda, A. Urano, J. Sakamoto, Y Imai, J Nucl. Mater. 258-263 (1998)159 interphase between the fibers and matrix, the fracture [9 C. Labrugere, L Guillaumat, A Guette, R. Naslain, J. Eur. Ceram Soc. 1 displacement increased with the increasing weight loss after (1991)641 oxidation and the fracture mode showed a non-brittle pattern. [10] H Ichikawa, Ann. Chim. Sci. Mater. 25(2000)523 [11M. Takeda, J. Sakamoto, Y Imai, H. Ichikawa, Compos. Sci. TechnoL. Acknowledgements (1999)813 [12]J F. Jamet, P. Lamicq, in: R. Naslain, J. Lamon, D. Doumeingts(eds ) High Temperature Ceramic Matrix Composites I, HT-CMC L, Woodhead The authors acknowledge the of the Publishing Limited, Cambridge, England, 1993, p. 735 National Foundation for Natura under [13] W. Yang. A Kohyama, T. Noda, et al., J. Nucl. Mater. 307-311(2002) No. 90405015 and the NSFC Disti Young under Contract No. 50425208. 2004 [14] C.E. Ramberg, G. Cruciani, K E. Spear, R E. Tressler, C F. Ramberg Jr, J. Am. Ceram. Soc. 79(1996)2089. 15 M. D. Thouless, O. Sbaizero, L.S. Sigl, A G. Evans, J. Am. Ceram Soc. 72 References [16] T. Mah, M.G. Mendiratta, A P. Katz, R. Ruh, K.S. Mazdiyasni, J.Am. 1]B N. Cox, F.W. Zok, Curr Opin. Solid State Mater. Sci. 1(1996)666. Ceram Soc. 68(1985)C27. [2]R. Naslain, Compos. Sci. TechnoL. 64(2004)155
composite oxidized at different temperature indicated the same regularity. The residual flexural strength of the composite was governed by the consumption of PyC interphase, and the oxidation had few effects on the strength of the composite. Above 1200 °C, the residual flexural strength of the composites was governed by the strength degradation of the Hi–Nicalon fiber. 3. Because oxidation took place along and weakened the interphase between the fibers and matrix, the fracture displacement increased with the increasing weight loss after oxidation and the fracture mode showed a non–brittle pattern. Acknowledgements The authors acknowledge the support of the Chinese National Foundation for Natural Sciences under Contract No.90405015 and the NSFC Distinguished Young Scholar under Contract No.50425208, 2004. References [1] B.N. Cox, F.W. Zok, Curr. Opin. Solid State Mater. Sci. 1 (1996) 666. [2] R. Naslain, Compos. Sci. Technol. 64 (2004) 155. [3] P. Fenici, A.J. Frias Rebelo, R.H. Jones, A. Kohyama, L.L. Snead, J. Nucl. Mater. 258–263 (1998) 215. [4] R. Naslain, Composites, Part A 29A (1998) 1145–1155. [5] S. Jacques, A. Lopez–Marure, C. Vincent, H. Vincent, J. Bouix, J. Eur. Ceram. Soc. 20 (2000) 1929. [6] R.J. Kerans, R.S. Hay, T.A. Parthasarathy, Curr. Opin. Solid State Mater. Sci. 4 (1999) 445. [7] R.E. Tressler, Composites, Part A 30 (1999) 429. [8] M. Takeda, A. Urano, J. Sakamoto, Y. Imai, J. Nucl. Mater. 258–263 (1998) 1594. [9] C. Labrugère, L. Guillaumat, A. Guette, R. Naslain, J. Eur. Ceram. Soc. 17 (1991) 641. [10] H. Ichikawa, Ann. Chim. Sci. Mater. 25 (2000) 523. [11] M. Takeda, J. Sakamoto, Y. Imai, H. Ichikawa, Compos. Sci. Technol. 59 (1999) 813. [12] J.F. Jamet, P. Lamicq, in: R. Naslain, J. Lamon, D. Doumeingts (eds.), High Temperature Ceramic Matrix Composites I, HT–CMC I, Woodhead Publishing Limited, Cambridge, England, 1993, p. 735. [13] W. Yang, A. Kohyama, T. Noda, et al., J. Nucl. Mater. 307–311 (2002) 1088. [14] C.E. Ramberg, G. Cruciani, K.E. Spear, R.E. Tressler, C.F. Ramberg Jr., J. Am. Ceram. Soc. 79 (1996) 2089. [15] M.D. Thouless, O. Sbaizero, L.S. Sigl, A.G. Evans, J. Am. Ceram. Soc. 72 (1989) 525. [16] T. Mah, M.G. Mendiratta, A.P. Katz, R. Ruh, K.S. Mazdiyasni, J. Am. Ceram. Soc. 68 (1985) C27. S. Wu et al. / Materials Letters 60 (2006) 3197–3201 3201