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/SicBelow 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 temper￾ature 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