obtained from Oak Ridge National Laboratory. The speci- and outer loading spans of 15 and 30 mm, respectively. The mens were fabricated by densifying multiple layers of two- flexure bar specimens were typically 3 X 4X 32 mm and dimensional plain-weave Nicalon fiber mats stacked in a they were loaded at a rate of 1. 27 mm min under ambient 0-30-60 layup sequence in a graphite die. Chemical conditions. The ultimate stress was determined from the vapor infiltration( CVI), under forced conditions of thermal peak load obtained from the load vs specimen displacement and pressure gradients, was used to densify the preforms plot. The WoF of the composites was estimated from the with Sic. The resulting composites were close to 90% area under the load vs specimen displacement plots, normal dense. Details of composite specimen fabrication are ized on the basis of unit cross-sectional area of the fractured described elsewhere [6,71 composites. The true specimen displacement was obtained The role of residual stresses in controlling the fiber/matrix by subtracting the system displacement from the total interfacial bond strength was evaluated by using unidirec- displacement. The system displacement was determined tional monofilament SiC-reinforced reaction-bonded silicon by measuring system compliance with the aid of a stiff nitride(rBsn) composite specimens that were fabricated at alumina piece the NaSa Lewis Research Center. Cleveland. Oh Fractured composite specimens were examined with a Commercially available uncoated (SCS-0)or carbon-coated scanning electron microscope(Model JXA-840A, JEOL (SCS-6) SiC monofilaments(Textron Specialty Materials Co, Ltd, Tokyo) to locate the failure origin and establish Division, Lowell, MA) were used as reinforcements. the associated characteristic fracture surface morphology of These monofilaments consist of a central carbon core. the fibers to estimate in-situ fiber strength 33 um in diameter, an intermediate layer of chemical vapor-deposited SiC and an optional outer 3 um thick carbon-rich double coating, resulting in an overall filament 3 Results and discussion diameter of s 142 um. Composites were fabricated by high-temperature consolidation and nitridation of alter- Fig. I shows the variation in measured fracture stress(i.e nately stacked SiC fiber mats and silicon cloth in a die. ultimate strength) with fiber coating thickness. The ultimate The details of the processing are described in Bhatt [81 strength increases with the coating thickness and reaches a Composite specimens with varying fiber content were peak value of 380 MPa at a coating thickness of =0.2- processed to change the interfacial residual stress. The 0.6 um. Further increase in coating thickness did not sig- effects of the changes in residual stress were estimated by nificantly affect the ultimate strength. A plot of WoF as a measuring interfacial bond strength using fiber pushout function of coating thickness, presented in Fig. 2, shows that tests WoF increases rapidly with fiber coating thickness and described by Bright et al. [9] Pushout specimens, / e reaches a peak value of 1.92 x 10*Nm m- at a coating The fiber pushout testing followed the procedur thickness of 0.13-0.3 um. The WoF did not change 4 mm thick. were cut normal to the fiber axis. The faces significantly with further increase in coating thickness normal to the fiber axis were ground flat and parallel to In-situ fiber strength of the composites was evaluated each other. Subsequently, one face was fine-polished from the characteristic markings on the fracture surface of 0. 25 um surface finish. These specimens were mounted on a the fibers in a Nicalon-fiber-reinforced SiC composite specimen holder with an =0.8 mm wide slot to allow the tested in the four-point-bend mode, as shown in Fig. 3 exit of fiber that is pushed out from the other end. The Characteristic features associated with brittle failure, such sample holder was fixed to an X-y micropositioning stage as mirror(smooth region around the fracture origin)and and the whole assembly was attached to a load cell mounted hackle( region of multiple fracture planes)are clearly obser- on the crosshead of a mechanical testing machine(Model vable on the surface of fractured fibers. SEM of the fibers 4505, Instron Corporation, Canton, MA). The pushout tests showed that most failed from defects or flaws on the fiber involved positioning of individual fibers(with the X-y surface. Using the measured values of mirror radii, we esti stage, as described in Bright et al. [9]) under an =60 um mated the tensile strength of the fibers from the following diameter flat tip tungsten carbide indenter fixed to the top empirical relationship, which is commonly used for cera- plate of the mechanical tester, followed by loading the fibers mics [10, Il] at a displacement rate of 0.05 mm min The load-displa- cement plots were recorded and the debonding and frictional shear stresses were calculated from the corresponding load where Im represents the mirror radius, or is the tensile and interface area. For each composite, 15-20 fibers were strength, and Am is the mirror constant, which is related to pushed out to account for the statistical variation in the local the fracture toughness of the material. In the present study. response of the materials Am is taken as 3.5 MPa m -, following the work of Thou Ultimate strength and WoF of Nicalon-fiber-reinforced less et al. [3] The strength of more than 30 Nicalon fibers for Sic matrix comp with fiber coating of varying thick- each set of composite specimens was determined by ness were obtained by fracturing composites in a four-point- measuring their fracture mirror radius and solving Eq. (1) bend mode on the mechanical testing machine with inner The strength distribution of fibers in the composites wasobtained from Oak Ridge National Laboratory. The speci￾mens were fabricated by densifying multiple layers of two￾dimensional plain-weave Nicalon fiber mats stacked in a 08–308–608 layup sequence in a graphite die. Chemical vapor infiltration (CVI), under forced conditions of thermal and pressure gradients, was used to densify the preforms with SiC. The resulting composites were close to 90% dense. Details of composite specimen fabrication are described elsewhere [6,7]. The role of residual stresses in controlling the fiber/matrix interfacial bond strength was evaluated by using unidirec￾tional monofilament SiC-reinforced reaction-bonded silicon nitride (RBSN) composite specimens that were fabricated at the NASA Lewis Research Center, Cleveland, OH. Commercially available uncoated (SCS-0) or carbon-coated (SCS-6) SiC monofilaments (Textron Specialty Materials Division, Lowell, MA) were used as reinforcements. These monofilaments consist of a central carbon core, 33 mm in diameter, an intermediate layer of chemical￾vapor-deposited SiC and an optional outer 3 mm thick carbon-rich double coating, resulting in an overall filament diameter of < 142 mm. Composites were fabricated by high-temperature consolidation and nitridation of alter￾nately stacked SiC fiber mats and silicon cloth in a die. The details of the processing are described in Bhatt [8]. Composite specimens with varying fiber content were processed to change the interfacial residual stress. The effects of the changes in residual stress were estimated by measuring interfacial bond strength using fiber pushout tests. The fiber pushout testing followed the procedure described by Bright et al. [9] Pushout specimens, < 1– 4 mm thick, were cut normal to the fiber axis. The faces normal to the fiber axis were ground flat and parallel to each other. Subsequently, one face was fine-polished to a 0.25 mm surface finish. These specimens were mounted on a specimen holder with an <0.8 mm wide slot to allow the exit of fiber that is pushed out from the other end. The sample holder was fixed to an X–Y micropositioning stage and the whole assembly was attached to a load cell mounted on the crosshead of a mechanical testing machine (Model 4505, Instron Corporation, Canton, MA). The pushout tests involved positioning of individual fibers (with the X–Y stage, as described in Bright et al. [9]) under an <60 mm diameter flat tip tungsten carbide indenter fixed to the top plate of the mechanical tester, followed by loading the fibers at a displacement rate of 0.05 mm min21 . The load–displa￾cement plots were recorded and the debonding and frictional shear stresses were calculated from the corresponding load and interface area. For each composite, 15–20 fibers were pushed out to account for the statistical variation in the local response of the materials. Ultimate strength and WOF of Nicalon-fiber-reinforced SiC matrix composites with fiber coating of varying thick￾ness were obtained by fracturing composites in a four-point￾bend mode on the mechanical testing machine with inner and outer loading spans of 15 and 30 mm, respectively. The flexure bar specimens were typically 3 × 4 × 32 mm and they were loaded at a rate of 1.27 mm min21 under ambient conditions. The ultimate stress was determined from the peak load obtained from the load vs specimen displacement plot. The WOF of the composites was estimated from the area under the load vs specimen displacement plots, normal￾ized on the basis of unit cross-sectional area of the fractured composites. The true specimen displacement was obtained by subtracting the system displacement from the total displacement. The system displacement was determined by measuring system compliance with the aid of a stiff alumina piece. Fractured composite specimens were examined with a scanning electron microscope (Model JXA-840A, JEOL Co., Ltd, Tokyo) to locate the failure origin and establish the associated characteristic fracture surface morphology of the fibers to estimate in-situ fiber strength. 3. Results and discussion Fig. 1 shows the variation in measured fracture stress (i.e. ultimate strength) with fiber coating thickness. The ultimate strength increases with the coating thickness and reaches a peak value of 380 MPa at a coating thickness of <0.2– 0.6 mm. Further increase in coating thickness did not sig￾nificantly affect the ultimate strength. A plot of WOF as a function of coating thickness, presented in Fig. 2, shows that WOF increases rapidly with fiber coating thickness and reaches a peak value of 1.92 × 104 Nm m22 at a coating thickness of <0.13–0.3 mm. The WOF did not change significantly with further increase in coating thickness. In-situ fiber strength of the composites was evaluated from the characteristic markings on the fracture surface of the fibers in a Nicalon-fiber-reinforced SiC composite, tested in the four-point-bend mode, as shown in Fig. 3. Characteristic features associated with brittle failure, such as mirror (smooth region around the fracture origin) and hackle (region of multiple fracture planes) are clearly obser￾vable on the surface of fractured fibers. SEM of the fibers showed that most failed from defects or flaws on the fiber surface. Using the measured values of mirror radii, we esti￾mated the tensile strength of the fibers from the following empirical relationship, which is commonly used for cera￾mics [10,11]: sfr 1=2 m ˆ Am …1† where rm represents the mirror radius, sf is the tensile strength, and Am is the mirror constant, which is related to the fracture toughness of the material. In the present study, Am is taken as 3.5 MPa m21/2, following the work of Thou￾less et al. [3] The strength of more than 30 Nicalon fibers for each set of composite specimens was determined by measuring their fracture mirror radius and solving Eq. (1). The strength distribution of fibers in the composites was 446 J.P. Singh et al. / Composites: Part A 30 (1999) 445–450