BRENNAN: INTERFACIAL CHARACTERIZATION Hi-Nicalon 5HS SyIramicTM 5HS 1200"C, 160 MPa, 19 cycles, 38 hrs to failure 1200.C, 160 MPa, 501 cycles, 1002 hrs to runout (Residual 1200C UTS=248 MPa, Ef=0.21%) Fig. 5. Fracture surface comparison between Hi-Nicalon and Sylramic SiC fiber MI SiC/SiC composites after 200° C tensile fatigue been found in many ceramic-matrix composites con- much less fibrous and with much shorter fiber pull taining the oxygen-rich Nicalon SiC fibers [14]. This out. However, the mode of fracture does not change carbon-rich layer is not seen in CVI SiC-matrix com- after elevated-temperature fatigue testing, nor does posites with Hi-Nicalon fibers and BN interfaces, the residual tensile strength change significantly which are processed at much lower temperatures TEM thin-foil analysis was performed on both as-fab- (1000oC) than the MI composites ricated and high-temperature-fatigued Sylramic fiber TEM thin-foil analysis of the Hi-Nicalon fiber MI composites, with the results indicating that no carbon composite that was subjected to tensile fatigue at rich layer forms during composite processing, nor 650C, as was shown in Fig. 6, and resulted in a quite does a glassy silica layer form as a result of high weak and brittle composite, found that a glassy silica temperature fatigue testing. This is illustrated in Fig layer had replaced the carbon-rich layer between the 9 for the Sylramic fiber composite sample that was BN and Hi-Nicalon fiber surface. This glassy silica subjected to 650C tensile fatigue testineoon-nich layer, as shown in Fig. 8, appears to be a result of shown previously in Fig. 6. The lack of a car oxidation of the carbon-rich layer and then the Hi- layer next to fiber surface is undoubtedly due to the Nicalon fiber surface due to matrix cracks forming higher temperature stability of the stoichiometric, during the fatigue testing. This glassy layer appar- crystalline SiC Sylramic fiber, when compared with ently bonds the Bn strongly to the fiber and may the carbon-rich Hi-Nicalon SiC fiber, During fatigue weaken the Hi-Nicalon fiber itself. At higher tempera- at elevated temperatures, any matrix cracks that may tures, such as that experienced during the 1200oC form do not cause pipeline oxidation down the fatigue test, this glassy layer can actually start to con- fiber/BN interface without the presence of the carbon sume the BN layer, totally bonding up the fiber/matrix rich layer interface, as was seen around the periphery of the At high stresses during high-temperature tensile 1200 C fatigue fracture surface shown in Fig. 5 fatigue, cracks do form in the Sylramic fiber MI com- As shown previously, in contrast to the fibrous fast posites, as shown for a sample in Fig. 10 that was fracture surface of Hi-Nicalon fiber MI composites, tensile fatigued at 815C at a stress of 186 MPa, the fracture surface of Sylramic fiber composites is which is well above the matrix microcracking stressBRENNAN: INTERFACIAL CHARACTERIZATION 4623 Fig. 5. Fracture surface comparison between Hi-Nicalon and Sylramic SiC fiber MI SiC/SiC composites after 1200°C tensile fatigue. been found in many ceramic-matrix composites con￾taining the oxygen-rich Nicalon SiC fibers [14]. This carbon-rich layer is not seen in CVI SiC-matrix com￾posites with Hi-Nicalon fibers and BN interfaces, which are processed at much lower temperatures (|1000°C) than the MI composites. TEM thin-foil analysis of the Hi-Nicalon fiber MI composite that was subjected to tensile fatigue at 650°C, as was shown in Fig. 6, and resulted in a quite weak and brittle composite, found that a glassy silica layer had replaced the carbon-rich layer between the BN and Hi-Nicalon fiber surface. This glassy silica layer, as shown in Fig. 8, appears to be a result of oxidation of the carbon-rich layer and then the Hi￾Nicalon fiber surface due to matrix cracks forming during the fatigue testing. This glassy layer appar￾ently bonds the BN strongly to the fiber and may weaken the Hi-Nicalon fiber itself. At higher tempera￾tures, such as that experienced during the 1200°C fatigue test, this glassy layer can actually start to con￾sume the BN layer, totally bonding up the fiber/matrix interface, as was seen around the periphery of the 1200°C fatigue fracture surface shown in Fig. 5. As shown previously, in contrast to the fibrous fast fracture surface of Hi-Nicalon fiber MI composites, the fracture surface of Sylramic fiber composites is much less fibrous and with much shorter fiber pull￾out. However, the mode of fracture does not change after elevated-temperature fatigue testing, nor does the residual tensile strength change significantly. TEM thin-foil analysis was performed on both as-fab￾ricated and high-temperature-fatigued Sylramic fiber composites, with the results indicating that no carbon￾rich layer forms during composite processing, nor does a glassy silica layer form as a result of high￾temperature fatigue testing. This is illustrated in Fig. 9 for the Sylramic fiber composite sample that was subjected to 650°C tensile fatigue testing, as was shown previously in Fig. 6. The lack of a carbon-rich layer next to fiber surface is undoubtedly due to the higher temperature stability of the stoichiometric, crystalline SiC Sylramic fiber, when compared with the carbon-rich Hi-Nicalon SiC fiber. During fatigue at elevated temperatures, any matrix cracks that may form do not cause pipeline oxidation down the fiber/BN interface without the presence of the carbon￾rich layer. At high stresses during high-temperature tensile fatigue, cracks do form in the Sylramic fiber MI com￾posites, as shown for a sample in Fig. 10 that was tensile fatigued at 815°C at a stress of 186 MPa, which is well above the matrix microcracking stress