《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_j.1551-2916.2005.00546.x[1]

J.Am. Ceram.Soc.,88p308-309502005 Dol:10.1l1551-29162005.00546x c 2005 The American Ceramic Society No claim to original US government works Journal Mechanical Properties of Thin Pyrolitic Carbon Interphase siC-Matrix Composites Reinforced with Near-Stoichiometric SiC Fibers Yutai Katoh t Takashi nozawa and Lance L. snead Metals and Ceramics Division, Oak Ridge National Laboratory, PO Box 2008, Oak Ridge, Tennessee 37831 Tensile properties of Tyranno-SA near-stoichiometric silicon spectively. Early generations of Sic/Sic composites produced carbide(SiC)-fiber-reinforced chemically vapor-infiltrated Sic before the mid-1990s were highly susceptible to neutron dama natrix composites with pyrolytic carbon interphases were ex- and exhibited severe strength degradation before acquiring a few rimentally studied. The influence of interphase thickness in a dpa of irradiation damage The strength reduction was attrib- ange of 60-300 nm on the tensile properties of the materials uted to interfacial debonding because of the significant irradia appeared to be generally minor. Thin interphase(600 nm)at the expense of composite stiffness. n research reactors to prove their largely improved irradiation According to irradiation studies reported for chemically vapor-deposited(CVD) SiC, it is very likely that polycrystalline B-SiC maintains a mechanical strength up to high neutron flue- ILICON CARBIDE(SiC) continuous fiber-reinforced SiC matrix ces at temperatures of interest for nuclear applications(roughly opposites(SiC/SiC composites) are promising structural 500-1000 C). 10, I This implies that irradiation effects on either erials for internal components of nuclear reactors. Partic- the cvi-SiC matrix or near-stoichiometric SiC fiber might not ularly, SiC/SiC composites are considered for the manufacture cause severe irradiation-induced degradation of the composites f control rod parts in very high temperature reactors (VHTRs) Therefore, the behavior of the PyC interphases during irradia- which is the concept of a gas-cooled thermal fission reactor with ion is most likely to determine the irradiation response of the ceramic fuel, based on the expectation that replacing carbon- composites. Generally, carbon materials are much more suscep- based materials with SiC/SiC composites confers a much longer tible to irradiation damage than SiC.It was reported that lifetime to the components and hence ultimately reduces the to on irradiation up to 10 dpa at 600C caused very significant tal cost of energy. SiC/SiC composites are also considered as microstructural modification in the Pyc interphase in near- very attractive materials for core components in gas-cooled fast stoichiometric SiC fiber composites, while the other constitu breeder reactors, blanket /first wall structures and inserts in fu- ents retained the microstructural stability. In another experi- sion power reactors, ,and intermediate heat exchangers of gas- ment, an obvious degradation of single-fiber push-in shear Doled nuclear systems' in the long term. Such high prospects strength of the fiber-matrix interface was observed for SiC/Sic composites in nuclear energy systems are based ichiometric SiC fiber-CVI-SiC compos th a 600-nm-thick various preferred properties of the B-phase Sic such as superior thermo-physical and thermo-mechanical properties, generally and 500 C 14 Therefore, one of the potential directions to fur- reasonable corrosion resistance, inherently low induced rad ther improve the irradiation tolerance of Sic/SiC composi activity and low decay heat, and excellent neutron irradiation to introduce significantly thinner PyC interphases than those tolerance usually applied In the present work, the influence of Pyc in e. The service environment of SiC/SiC composites for nuclear terphase thickness on tensile fast fracture properties in near- plications involves a strong neutron radiation field. while oth- stoichiometric SiC fiber-CVI-SiC matrix composites was studied er conditions like temperature and oxidative environment will be systematically relatively mild. For example, control rods and guide tubes of VhTR are expected to acquire a lifetime fast neutron dose of 2x 10-n/m(E>0.1 MeV), which corresponds to a damage Il. Experimental Procedure level of approximately 20 displacements per atom(dpa)in SiC. hile the helium coolant in VhTR pressure vessel will have a The composites used in this study were produced through a normal outlet temperature and oxygen /water vapor concentra forced-flow thermal gradient chemical vapor infiltration (F- tions of 1000C and s0.1 ppm(for both O, and H,O) CVi) process developed at Oak Ridge National Laboratory The reinforcement consisted of Tyranno-SA Grade-3(SA3)- sintered SiC fibers, with an average diameter of < 7.5 um, in a T.A. Parthasarathy--contributing editor 2D plain-woven architecture(fiber and cloth lot numbers Psa SI7116PX and SA3-S1ll6Px, respectively). The properties of Tyranno-SA3 fiber are summarized in Table I. The thread cript No 20445. Received September 28, 2004: approved April 22, 2005 count of the fabric was 17 yarns/in., and the number of filaments elle LLC, "JUPIT per yarn was 1600 nominal. For each composite, approximately cience and Technology 70 fabrics punched into 3-in.-diameter circles were stacked in ation([0/90] hereafter)and held tightly in a tructural Materials in Mixed-Spectrum Fission Reactors. graphite fixture for interphase deposition and subsequent matrix Author to whom correspondence should be addressed. e-mail: kathy oml go densification

Mechanical Properties of Thin Pyrolitic Carbon Interphase SiC–Matrix Composites Reinforced with Near-Stoichiometric SiC Fibers Yutai Katoh,w Takashi Nozawa, and Lance L. Snead Metals and Ceramics Division, Oak Ridge National Laboratory, PO Box 2008, Oak Ridge, Tennessee 37831 Tensile properties of Tyrannot-SA near-stoichiometric silicon carbide (SiC)-fiber–reinforced chemically vapor-infiltrated SiC￾matrix composites with pyrolytic carbon interphases were ex￾perimentally studied. The influence of interphase thickness in a range of 60–300 nm on the tensile properties of the materials appeared to be generally minor. Thin interphase (o100 nm) did not have a significant deteriorating effect on composite proper￾ties, which has commonly been reported for conventional SiC- fiber composites. For very thin interphase (o60 nm) composites, a slight decrease in fracture strain and a substantial increase in interfacial sliding stress were noted. Increases in ultimate tensile strength and fracture strain were observed at a much thicker interphase (4600 nm) at the expense of composite stiffness. I. Introduction SILICON CARBIDE (SiC) continuous fiber–reinforced SiC matrix composites (SiC/SiC composites) are promising structural materials for internal components of nuclear reactors.1 Partic￾ularly, SiC/SiC composites are considered for the manufacture of control rod parts in very high temperature reactors (VHTRs), which is the concept of a gas-cooled thermal fission reactor with ceramic fuel, based on the expectation that replacing carbon￾based materials with SiC/SiC composites confers a much longer lifetime to the components and hence ultimately reduces the to￾tal cost of energy. SiC/SiC composites are also considered as very attractive materials for core components in gas-cooled fast breeder reactors,2 blanket/first wall structures and inserts in fu￾sion power reactors,3,4 and intermediate heat exchangers of gas￾cooled nuclear systems5 in the long term. Such high prospects for SiC/SiC composites in nuclear energy systems are based on various preferred properties of the b-phase SiC such as superior thermo-physical and thermo-mechanical properties, generally reasonable corrosion resistance, inherently low induced radio￾activity and low decay heat, and excellent neutron irradiation tolerance.6 The service environment of SiC/SiC composites for nuclear applications involves a strong neutron radiation field, while oth￾er conditions like temperature and oxidative environment will be relatively mild. For example, control rods and guide tubes of VHTR are expected to acquire a lifetime fast neutron dose of 42
1026 n/m2 (E40.1 MeV), which corresponds to a damage level of approximately 20 displacements per atom (dpa) in SiC, while the helium coolant in VHTR pressure vessel will have a normal outlet temperature and oxygen/water vapor concentra￾tions of B10001C and B0.1 ppm (for both O2 and H2O), re￾spectively. Early generations of SiC/SiC composites produced before the mid-1990s were highly susceptible to neutron damage and exhibited severe strength degradation before acquiring a few dpa of irradiation damage.7 The strength reduction was attrib￾uted to interfacial debonding because of the significant irradia￾tion-induced shrinkage of non-stoichiometric SiC(-based) fibers (e.g., ceramic-grade [CG] Nicalont and Hi-Nicalont, Nippon Q2 Carbon Co., Ltd., Tokyo, Japan). As near-stoichiometric and high-crystallinity SiC fibers (Hi-Nicalont Type-S and Tyran￾not-SA, Ube Industries, Ltd., Tokyo, Japan) became available, Q3 composites with these fibers, pyrolytic carbon (PyC) interphase, and chemically vapor infiltrated (CVI) SiC matrices were placed in research reactors to prove their largely improved irradiation resistance.8,9 According to irradiation studies reported for chemically vapor-deposited (CVD) SiC, it is very likely that polycrystalline b-SiC maintains a mechanical strength up to high neutron fluen￾ces at temperatures of interest for nuclear applications (roughly 5001–10001C).10,11 This implies that irradiation effects on either the CVI-SiC matrix or near-stoichiometric SiC fiber might not cause severe irradiation-induced degradation of the composites. Therefore, the behavior of the PyC interphases during irradia￾tion is most likely to determine the irradiation response of the composites. Generally, carbon materials are much more suscep￾tible to irradiation damage than SiC.12 It was reported that ion irradiation up to 10 dpa at 6001C caused very significant microstructural modification in the PyC interphase in near￾stoichiometric SiC fiber composites, while the other constitu￾ents retained the microstructural stability.13 In another experi￾ment, an obvious degradation of single-fiber push-in shear strength of the fiber–matrix interface was observed in near-sto￾ichiometric SiC fiber–CVI-SiC composites with a 600-nm-thick PyC interphase after neutron irradiation up to 0.5 dpa at 3001 and 5001C.14 Therefore, one of the potential directions to fur￾ther improve the irradiation tolerance of SiC/SiC composites is to introduce significantly thinner PyC interphases than those usually applied. In the present work, the influence of PyC in￾terphase thickness on tensile fast fracture properties in near￾stoichiometric SiC fiber–CVI-SiC matrix composites was studied systematically. II. Experimental Procedure The composites used in this study were produced through a forced-flow thermal gradient chemical vapor infiltration (F￾CVI) process developed at Oak Ridge National Laboratory.15 The reinforcement consisted of Tyrannot-SA Grade-3 (SA3)- sintered SiC fibers, with an average diameter of B7.5 mm, in a 2D plain-woven architecture (fiber and cloth lot numbers PSA￾S17I16PX and SA3-S1I16PX, respectively).16 The properties of Tyrannot-SA3 fiber are summarized in Table I. The thread count of the fabric was 17 yarns/in., and the number of filaments per yarn was 1600 nominal. For each composite, approximately 70 fabrics punched into 3-in.-diameter circles were stacked in a [01/901] orientation ([0/90] hereafter) and held tightly in a graphite fixture for interphase deposition and subsequent matrix densification. Journal J. Am. Ceram. Soc., 88 [11] 3088–3095 (2005) DOI: 10.1111/j.1551-2916.2005.00546.x r 2005 The American Ceramic Society No claim to original US government works 3088 T. A. Parthasarathy—contributing editor Supported by the Office of Fusion Energy Sciences, U.S. Department of Energy under contract DE-AC05-00OR22725 with UT-Battelle, LLC, ‘‘JUPITER-II’’ U.S.-Department of Energy/Japanese Ministry of Education, Culture, Sports, Science and Technology (MEXT) collaboration for fusion material system research, and U.S.-Department of Ener￾gy/ Japan Atomic Energy Research Institute (JAERI) Collaborative Program on FWB Structural Materials in Mixed-Spectrum Fission Reactors. w Author to whom correspondence should be addressed. e-mail: katohy@ornl.gov Manuscript No. 20445. Received September 28, 2004; approved April 22, 2005

ovember 2005 Mechanical Properties of Thin Py C Interphase Sic-Matrix Composites 3089 Table I. Properties of a Tyranno"-SA Grade-3 Fiber the range of 26-54 nm. as shown in Table il. the matrix densit was high in the center regardless of the vertical position and was Tyranno"-SA Grade-3 lowest at the botton ge. The(micro-) structural uniformity Atomic composition SiCLoSAlooos Diameter(um) Number of filaments/yarn A miniature tensile specimen geometry that had been devel- oped for neutron irradiation studies was used for the tensile Tensile strength(GPa) testing. The gauge dimensions were as follows: 15-mm length Tensile modulus(GPa) 400 3-mm width and <2.3-mm thickness. The gauge width was se- Mass density (g/cm) 3.l lected such that each fabric layer was able to accommodate two fiber strands, which is the minimum requirement to avoid the potentially significant gauge width effect imposed by an insuffi- The PyC interphase was deposited in an isothermal configura- cient number of fiber strands. The gauge thickness accommo- tion using propylene as the precursor at 1100C, a total pressure dated approximately 15 fabric layers. The testing was performed of 5 kPa, and flow rates of 50 and 1000 cm/min for propylea following the general guidelines of AsTM C1275 and C1359 at a and diluting argon, respectively. The deposition rate of PyC was crosshead displacement rate of 0.5 mm/min. For the ambient I nm/min. The thickness of the interphase was controlled by temperature tests, the specimens were clamped on both faces by wedge grips with aluminum tabs between the specimen and using methyltrichlorosilane (MTS, Gelest Inc, Tullytown, Pa) grips, and the strain was measured using a set of strain gauges at the fixture's hot surface temperature of 1050oC and a back attached on both faces. The magnitude of the bending strain pressure of 100 kPa. The mTs precursor was carried by hy component appeared to be less than 10% of the tensile strain in drogen bubbling at flow rates for MTS and the carrier of 0.3 g/ most cases. The mean strain was used for analysis min and 450 cm/ min, respectively. The deposition rate was es- The elevated temperature tensile test was conducted at timated to be 20 nm/min. a thin Sic coating was applied to 1300.C in a commercial argon fow. The estimated oxygen par some of the composites prior to the Pyc interphase deposition in tial pressure during heating and testing was c0. I Pa. The spec- order to modify the effective surface roughness of the fibers. In imens were loaded after I h of equilibration period at the test such cases, the coating was deposited at one-half of the flow temperature, at which the entire specimen was maintained dur- rates for both MTS and hydrogen and at 1100.C in an isother ng the testing. Only limited test runs were equipped with laser mal configuration. A list of materials is provided in Table Il. extensometry, while most of the testings were performed without The variation of Pyc interphase thickness within the 3-in. di a dedicated strain measurement at the specimens. Details of the elevated temperature testing are reported elsewhere Tensile specimens were machined out of the infiltrated com- posites in a [0/90 orientation. Small rectangular blocks were also cut from a location right next to each tensile specimen for microstructural characterization and to allow measurement of the interphase thickness from polished surfaces using a field (1) Microstructural examinatio emission scanning electron microscope. Standard deviations for Topographical roughness of the fiber-interphase interface was the interphase thickness within individual blocks were < 10% characterized by a typical peak-to-peak amplitude of 20 nm Selected sample blocks were also examined by transmission and an average root-mean-square(rMs)deviation of 3. 7 nm. electron microscopy for closer examination of the interfacial This corresponds to the features of faceted blocky SiC crystal microstructures and the PyC interphase Fiber volume fraction grains, the sizes of which appeared to be 30-150 nm as observed ind apparent mass density were determined using the individual by TEM, close to the surface of the Tyranno-SA3 fibers. A specimens. Generally, the Pyc interphase thickness varied to TEM image representing the interfacial structure in the single some extent within a 3-in. -diameter composite disk; relatively Pyc interphase composites is presented in Fig. 1. The high-res- thick at the bottom center(the reactant comes from below) and olution lattice images shown in Fig. 2 imply clean interfaces thin on the top. The relative deviation of the average Pyc in- both between B-SiC in the fiber and carbon interphase and be- terphase thickness for individual specimens was-40% to +30% tween the carbon interphase and B-SiC matrix. Also, electron of the average in a single composite disk; for example, the a energy loss spectra did not show any evidence of the presence of erage PyC interphase thickness in individual specimens of the oxygen or other impurities at the interface. The structure of the Table Il. List of Composites Material ID Cv-1266 CVI-1264 CvL-1265 CvI-1268 Cv-1269 CvI-1271 Reinforcement Fiber Tyranno-SA Grade-3(7. 5 um) Architecture 2D-Plain Weave [0/90] lay-up Fiber volume fraction (% 35.2 354 8.8 399 Interphase thickness(nm) Measured average one None None Range of scatte 0-74 73-l Measured average l16 Range of scatte 86-155 56-126101-141 588833 Matrix/composite 2.62 2.61 72 2.74 2.71 2.42 17.7 17.7 13.4 l5.1 14.4 23.5 tic carbon; CVI, chemically vapor infiltrated

The PyC interphase was deposited in an isothermal configura￾tion using propylene as the precursor at 11001C, a total pressure of 5 kPa, and flow rates of 50 and 1000 cm3 /min for propylene and diluting argon, respectively. The deposition rate of PyC was B1 nm/min. The thickness of the interphase was controlled by the time of deposition. The matrix infiltration was carried out using methyltrichlorosilane (MTS, Gelest Inc., Tullytown, PA) at the fixture’s hot surface temperature of 10501C and a back pressure of B100 kPa. The MTS precursor was carried by hy￾drogen bubbling at flow rates for MTS and the carrier of 0.3 g/ min and 450 cm3 /min, respectively. The deposition rate was es￾timated to be B20 nm/min. A thin SiC coating was applied to some of the composites prior to the PyC interphase deposition in order to modify the effective surface roughness of the fibers. In such cases, the coating was deposited at one-half of the flow rates for both MTS and hydrogen and at 11001C in an isother￾mal configuration. A list of materials is provided in Table II. The variation of PyC interphase thickness within the 3-in. di￾ameter disks was typically 730%. Tensile specimens were machined out of the infiltrated com￾posites in a [0/90] orientation. Small rectangular blocks were also cut from a location right next to each tensile specimen for microstructural characterization, and to allow measurement of the interphase thickness from polished surfaces using a field emission scanning electron microscope. Standard deviations for the interphase thickness within individual blocks were B10%. Selected sample blocks were also examined by transmission electron microscopy for closer examination of the interfacial microstructures and the PyC interphase. Fiber volume fraction and apparent mass density were determined using the individual specimens. Generally, the PyC interphase thickness varied to some extent within a 3-in.-diameter composite disk; relatively thick at the bottom center (the reactant comes from below) and thin on the top. The relative deviation of the average PyC in￾terphase thickness for individual specimens was –40% to 130% of the average in a single composite disk; for example, the av￾erage PyC interphase thickness in individual specimens of the composite CVI-1266 (average PyC thickness of 42 nm) varied in the range of 26–54 nm, as shown in Table II. The matrix density was high in the center regardless of the vertical position and was lowest at the bottom edge. The (micro-) structural uniformity for this particular F-CVI configuration has been studied in a companion work.17 A miniature tensile specimen geometry that had been devel￾oped for neutron irradiation studies was used for the tensile testing.18 The gauge dimensions were as follows: 15-mm length, 3-mm width and B2.3-mm thickness. The gauge width was se￾lected such that each fabric layer was able to accommodate two fiber strands, which is the minimum requirement to avoid the potentially significant gauge width effect imposed by an insuffi- cient number of fiber strands.19 The gauge thickness accommo￾dated approximately 15 fabric layers. The testing was performed following the general guidelines of ASTM C1275 and C1359 at a crosshead displacement rate of 0.5 mm/min. For the ambient temperature tests, the specimens were clamped on both faces by wedge grips with aluminum tabs between the specimen and grips, and the strain was measured using a set of strain gauges attached on both faces. The magnitude of the bending strain component appeared to be less than 10% of the tensile strain in most cases. The mean strain was used for analysis. The elevated temperature tensile test was conducted at 13001C in a commercial argon flow. The estimated oxygen par￾tial pressure during heating and testing was B0.1 Pa. The spec￾imens were loaded after 1 h of equilibration period at the test temperature, at which the entire specimen was maintained dur￾ing the testing. Only limited test runs were equipped with laser extensometry, while most of the testings were performed without a dedicated strain measurement at the specimens. Details of the elevated temperature testing are reported elsewhere.20 III. Results (1) Microstructural Examination Topographical roughness of the fiber–interphase interface was characterized by a typical peak-to-peak amplitude of B20 nm and an average root-mean-square (RMS) deviation of 3.7 nm. This corresponds to the features of faceted blocky SiC crystal grains, the sizes of which appeared to be 30–150 nm as observed by TEM, close to the surface of the Tyrannot-SA3 fibers. A TEM image representing the interfacial structure in the single PyC interphase composites is presented in Fig. 1. The high-res￾olution lattice images shown in Fig. 2 imply clean interfaces both between b-SiC in the fiber and carbon interphase and be￾tween the carbon interphase and b-SiC matrix. Also, electron energy loss spectra did not show any evidence of the presence of oxygen or other impurities at the interface. The structure of the PyC interphase became more like a glassy carbon closer to the Table I. Properties of a Tyrannot-SA Grade-3 Fiber Properties Tyrannot-SA Grade-3 Atomic composition SiC1.08Al0.005 Diameter (mm) 7.5 Number of filaments/yarn 1600 Tensile strength (GPa) 2.5 Tensile modulus (GPa) 400 Mass density (g/cm3 ) 3.1 Table II. List of Composites Material ID CVI-1266 CVI-1264 CVI-1265 CVI-1267 CVI-1268 CVI-1269 CVI-1271 Reinforcement Fiber Tyrannot-SA Grade-3 (7.5 mm) Architecture 2D-Plain Weave [01/901] lay-up Fiber volume fraction (%) 35.2 35.4 35.3 38.8 38.8 38.8 39.9 Interphase thickness (nm) SiC Measured average None None None 39 87 65 72 Range of scatter 0–74 73–111 58–75 56–80 PyC Measured average 42 116 226 93 120 207 648 Range of scatter 26–54 86–155 168–256 56–126 101–141 169–256 588–833 Matrix/composite Density (g/cm3 ) 2.62 2.61 2.72 2.74 2.69 2.71 2.42 Porosity (%) 17.4 17.7 17.7 13.4 15.1 14.4 23.5 PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Q1 November 2005 Mechanical Properties of Thin PyC Interphase SiC–Matrix Composites 3089

Journal of the American Ceramic Society--Katoh Central gauge section Peripheral gauge section 25 Shoulder section Ty InoSA t(=23) 13 millimeter Fig 1. Low-magnification TEM image showing the typical topograph ical roughness of a fiber-interphase interface ripheral gauge section), and 20% failed in the shoulder section very close to the gauge section. In Fig. 3, each section is illus rated along with the dimensions of the specimen. The central Sic matrix. Within the Pyc interphase, spondin ally dominating, while sp- bonding was dominant fiber layer. Therefore, the Pyc interphase is me near the fibers, and there is a gradual transition to 号 auge section corresponds approximately to the section to which strain gauges are attached. For the elevated temperature tests at 300C, failure probabilities at the central gauge section, the peripheral gauge section, and the shoulder section were approx along the radial direction mately 40%, 40%, and 20%, respectively. The reason for the increased failure probability at the peripheral gauge section at the elevated temperature is not known. Tests in which failure (2) Tensile Properties occurred within either of the gauge sections were considered Approximately 50%o of the specimens tested at ambient alid, because the failure probabilities for the central and pe- ature failed in a plane within +5 mm from the center ripheral gauge sections in the ambient temperature tests were gauge section) of the 15-mm-long gauge section, 30% approximately proportional to the section length. Representa- within 2.5 mm from the gauge ends in the gauge section(pe- Live stress-strain curves incorporating multiple unloading-re- Fig 4. All the specimens exhibited the initial proportional stage and the subsequent transition to the second linear stage that continues until the failure before apparent achievement of ma- trix crack saturation The results from the tensile tests are summarized in Table ill The tangential modulus was determined by the slope of SA fi straight reloading segment after unloading from 25 MPa of ten- sile stress. The proportional limit stress(PLS)was defined as the stress at 5% stress deviation from the extrapolated linear seg- ment used for the modulus determination This definition avoids 0.05% strain offset method specified in ASTM C1275. As seen in Fig. 4, the 0.05% strain offset lines would intersect the stress strain curves at stresses higher than 150 MPa. whereas deviation RNL Tyranno-SA/PyC/CV-SIC lac=250nm tac"650nm Tensile Strain(%) 2. olution TEM images of the interface region between an Fig. oSA Sic fiber interphase(A)and between a PyC interphase and Sic ive 0/90/PyC/CVI-SiC co [090J.[0°90PyC, matrix (B bon: CV, chemically vapor

SiC matrix. Within the PyC interphase, sp3 bonding was gener￾ally dominating, while sp2 bonding was dominant in the near- fiber layer.21 Therefore, the PyC interphase is more graphitic near the fibers, and there is a gradual transition to glassy carbon along the radial direction. (2) Tensile Properties Approximately 50% of the specimens tested at ambient temper￾ature failed in a plane within 75 mm from the center (central gauge section) of the 15-mm-long gauge section, B30% failed within 2.5 mm from the gauge ends in the gauge section (pe￾ripheral gauge section), and B20% failed in the shoulder section very close to the gauge section. In Fig. 3, each section is illus￾trated along with the dimensions of the specimen. The central gauge section corresponds approximately to the section to which strain gauges are attached. For the elevated temperature tests at 13001C, failure probabilities at the central gauge section, the peripheral gauge section, and the shoulder section were approx￾imately 40%, 40%, and 20%, respectively. The reason for the increased failure probability at the peripheral gauge section at the elevated temperature is not known. Tests in which failure occurred within either of the gauge sections were considered valid, because the failure probabilities for the central and pe￾ripheral gauge sections in the ambient temperature tests were approximately proportional to the section length. Representa￾tive stress–strain curves incorporating multiple unloading–re￾loading sequences obtained from the valid tests are presented in Fig. 4. All the specimens exhibited the initial proportional stage and the subsequent transition to the second linear stage that continues until the failure before apparent achievement of ma￾trix crack saturation. The results from the tensile tests are summarized in Table III. The tangential modulus was determined by the slope of a straight reloading segment after unloading from 25 MPa of ten￾sile stress. The proportional limit stress (PLS) was defined as the stress at 5% stress deviation from the extrapolated linear seg￾ment used for the modulus determination. This definition avoids unreasonable overestimation of PLS that occurs by using the 0.05% strain offset method specified in ASTM C1275. As seen in Fig. 4, the 0.05% strain offset lines would intersect the stress– strain curves at stresses higher than 150 MPa, whereas deviation Fig. 2. High-resolution TEM images of the interface region between an SiC fiber and PyC interphase (A) and between a PyC interphase and SiC matrix (B). PyC, pyrolytic carbon. Fig. 3. Geometry of miniature tensile specimen. Unit of dimensions is millimeter. Fig. 4. Representative load–strain curves obtained for Tyrannot-SA 2D Plain-Weave [0/90]/PyC/CVI-SiC composites. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Fig. 1. Low-magnification TEM image showing the typical topograph￾ical roughness of a fiber–interphase interface. 3090 Journal of the American Ceramic Society—Katoh et al. Vol. 88, No. 11

ovember 2005 Mechanical Properties of Thin PyC Interphase SiC-Matrix Composites Table Ill. Summary of Tensile Properties CVI-1266 CvI-1264 CvI-1265 CVI-1268 CVI-1269 CvI-1271 Interphase thickness(nm) Sic (measured average) PyC(measured average) Fiber volume fraction (% 35.2 38.8 38.8 Porosity (%) 174 17.7 15.1 144 Ambient Number of valid tests UTS(MPa) 2l1(201)233(33232(34)227(35) 34(17 24l(38) PLS (MPa) 61(17) 53(13) 62(14) 58(16) 45(12) 53(4) Modulus(GPa) 252(36) 192(30)213(23)228(19) 08(17)204(32) At 1573 K in aff,d tests Number of vali 0 UTS (MPa) 182(52)193(28)209(33)178(8) 233(18 PLS (MPa) 53(7) arentheses show standard deviations. n/m, not measured; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated: UTS, ultimate tensile stress: PLS, proportional limit stress. from the initial linear segments occurs at much lower stress lev multiplied by 1.5 to allow comparison, assuming that only the els. The ambient temperature ultimate tensile stress (UTS) and longitudinal fibers carry the load just before failure. As a general PLS were generally in ranges of 210-240 and 45-60 MPa,re- trend, the UTS was insensitive to tpyc in the range of 40-250 nm spectively, for all materials except CVl-1271, in which an ex- Composites with a PyC layer thinner than 40 nm appeared to ceptionally thick Pyc interphase (650 nm in average) had have a slightly lower UTS, although no statistical significan been applied. The tangential moduli were 190-250 GPa except could be determined for these differences. Composites with for CV1-1271, again implying negative dependence on the Pyc Pyc layer thicker than -600 nm exhibited a significantly larger layer thickness(Pvc). The additional SiC layer between the fiber UTS. Data points from other works2 fall approximately and the Pyc layer did not impose any noticeable effect on the within the data band associated with the materials investigated tensile properties of the composite. UTS at 1300.C in argon this work. The PLs and the modulus exhibited slightly neg- appeared to be 10%0-25% lower than at ambient temperature ative correlation with Ipyc in a similar manner. As a result, the Limited data at the elevated temperature with strain measure- proportional limit strain is nearly independent of Pvc, except at ment suggest that the tangential modulus decreases at 1300C by a very large thickness. The trend lines in Figs. 5 and 6 represent fractions similar to the decrease in UTS. However, it was not linear fits to the data points, whereas that in Fig. 7 represents a possible to determine the effect of testing temperature on PLs. fit by a simple model discussed later. The trend line in Fig. 5 was UTS, PLS, and tangential modulus measured for each tensile not extended to >300 nm because of the inability to determine pecimen are plotted against tpyc for individual specimens the shape of the curve that fits to the significantly higher UTS at Figs. 5-7, respectively. In 5, UTS data for similar compos- Ipyc 600 nm ites are plotted together. The Oak Ridge National Labora- tory(ORNL) Tyranno-SA/PyC/CVI-SiC [0/30/60] composites had been reinforced by fibers from the same production lot as (3) Examination of fractured specimens used in this work, while the National Institute for Materials The composites'fracture surfaces consist of matrices, longitu- Science(NIMS, Tsukuba, Japan) Tyranno-SA/PyC/CVI-Sic dinal fiber bundles, and transverse fiber bundles. Fracture sur [o/90]composites used Grade-3 fibers from a different lot. The faces of the longitudinal fiber bundles are characterized by fiber tensile strength of composites with a [0/30/60] fabric lay-up was pull-out with a length of 10-50 Hm. No significant difference 120 B ORNL Tyranno-SAPyC/CVISC D/9a Tyranno-SASICAPyCNCVI-SIC A ORNL Tyranc-SA/SiC/PyC/CV-SIC C/D NIMS Tyanno-SAPyckVISiIC 090 a NIMS Tyranno-SAPyC/CVI-SIC 0S A ORNL Tyranno-SAPyCACVISIC 0//6D(x1 5) x ORNL Hi-Nicalon- S/PyC/CV-SC 0/0/60(X1, 5% Na-nterphase PyC Interphase Thickness(nm) Fig. 5. Ultimate tensile stress of Tyranno-SA 2D Plain-Weave [o/90V (Sic/ Pyc/ CVI-SiC composites at ambient temperature plotted against PyC interphase thickness. 0/900/90] PyC, pyrolytic carbon; CVI slotted hemically vapor infiltrated

from the initial linear segments occurs at much lower stress lev￾els. The ambient temperature ultimate tensile stress (UTS) and PLS were generally in ranges of 210–240 and 45–60 MPa, re￾spectively, for all materials except CVI-1271, in which an ex￾ceptionally thick PyC interphase (B650 nm in average) had been applied. The tangential moduli were 190–250 GPa except for CVI-1271, again implying negative dependence on the PyC layer thickness (tPyC). The additional SiC layer between the fiber and the PyC layer did not impose any noticeable effect on the tensile properties of the composite. UTS at 13001C in argon appeared to be 10%–25% lower than at ambient temperature. Limited data at the elevated temperature with strain measure￾ment suggest that the tangential modulus decreases at 13001C by fractions similar to the decrease in UTS. However, it was not possible to determine the effect of testing temperature on PLS. UTS, PLS, and tangential modulus measured for each tensile specimen are plotted against tPyC for individual specimens in Figs. 5–7, respectively. In Fig. 5, UTS data for similar compos￾ites are plotted together.17,22 The Oak Ridge National Labora￾tory (ORNL) Tyrannot-SA/PyC/CVI-SiC [0/30/60] composites had been reinforced by fibers from the same production lot as used in this work, while the National Institute for Materials Science (NIMS, Tsukuba, Japan) Tyrannot-SA/PyC/CVI-SiC [0/90] composites used Grade-3 fibers from a different lot. The tensile strength of composites with a [0/30/60] fabric lay-up was multiplied by 1.5 to allow comparison, assuming that only the longitudinal fibers carry the load just before failure. As a general trend, the UTS was insensitive to tPyC in the range of 40–250 nm. Composites with a PyC layer thinner than B40 nm appeared to have a slightly lower UTS, although no statistical significance could be determined for these differences. Composites with a PyC layer thicker than B600 nm exhibited a significantly larger UTS. Data points from other works17,22 fall approximately within the data band associated with the materials investigated in this work. The PLS and the modulus exhibited slightly neg￾ative correlation with tPyC in a similar manner. As a result, the proportional limit strain is nearly independent of tPyC, except at a very large thickness. The trend lines in Figs. 5 and 6 represent linear fits to the data points, whereas that in Fig. 7 represents a fit by a simple model discussed later. The trend line in Fig. 5 was not extended to 4300 nm because of the inability to determine the shape of the curve that fits to the significantly higher UTS at tPyC\600 nm. (3) Examination of Fractured Specimens The composites’ fracture surfaces consist of matrices, longitu￾dinal fiber bundles, and transverse fiber bundles. Fracture sur￾faces of the longitudinal fiber bundles are characterized by fiber pull-out with a length of 10–50 mm. No significant difference Fig. 5. Ultimate tensile stress of Tyrannot-SA 2D Plain-Weave [0/90]/ (SiC/) PyC/ CVI-SiC composites at ambient temperature plotted against PyC interphase thickness. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Fig. 6. Proportional limit tensile stress of Tyrannot-SA 2D Plain￾Weave [01/901]/(SiC/) PyC/CVI-SiC composites at ambient temperature plotted against PyC interphase thickness. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Table III. Summary of Tensile Properties Material ID CVI-1266 CVI-1264 CVI-1265 CVI-1267 CVI-1268 CVI-1269 CVI-1271 Interphase thickness (nm) SiC (measured average) None None None 39 87 65 72 PyC (measured average) 42 116 226 93 120 207 648 Fiber volume fraction (%) 35.2 35.4 35.3 38.8 38.8 38.8 39.9 Porosity (%) 17.4 17.7 17.7 13.4 15.1 14.4 23.5 Ambient Number of valid tests 4 5 4 7 4 4 4 UTS (MPa) 211 (20w ) 233 (33) 232 (34) 227 (35) 234 (17) 241 (38) 304 (16) PLS (MPa) 61 (17) 53 (13) 62 (14) 58 (16) 45 (12) 53 (4) 46 (1) Modulus (GPa) 252 (36) 192 (30) 213 (23) 228 (19) 208 (17) 204 (32) 157 (5) At 1573 K in argon Number of valid tests 3 5 8 4 1 0 4 UTS (MPa) 182 (52) 193 (28) 209 (33) 178 (8) 177 n/m 233 (18) PLS (MPa) n/m n/m 49 (4) n/m n/m n/m 53 (7) Modulus (GPa) n/m n/m 199 (9) n/m n/m n/m 129 (30) w Numbers in parentheses show standard deviations. n/m, not measured; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated; UTS, ultimate tensile stress; PLS, proportional limit stress. November 2005 Mechanical Properties of Thin PyC Interphase SiC–Matrix Composites Q1 3091

3092 Journal of the American Ceramic Society-Katoh et al. Vol. 88. No. I 293K ORNL Tyranno-SA/AyCAVH-SIC C90 (a A ORNL TyrannD-SA/PyCKCVH-SIC 0/060 x ORNL Hi-Nicalon-S/Pyc/CVI-SC 0/3060 a"-1 No- intara Fig. 7. Influence of Pyc . Weave [0"/90y(Sicn) PyC/ CVI (B) SiC composites [0/90 [0 pyrolytic carbon; CV. chemically vapor infiltrated. was noticed for the pull-out length among com a Pvc of 25-250 nm, as compared in Figs. 8(A)a composite with a Pvc of 830 nm exhibited subs fiber pull-out lengths of 50-200 um as shown in Fig 8(C). De- bonding of the longitudinal fiber interfaces occurred mostly at the interface between the Pyc interphase and the matrix. The fiber fracture surfaces were dominated by grain boundary frac ture accompanied by river patterns, which is typical for Tyr- anno-SA fibers. Most of the fiber fractures originated at the urface (insets to Fig 8) but some originated within the core egion. Side surfaces of the pull-out fibers appeared free from matches or any evidence of wear damage, exhibiting only the original roughness because of the surface grain structures. As for the fracture surfaces of transverse fiber bundles, major cracking took place mostly within the fiber bundles, and the debond wa observed at both the fiber-PyC and the Pyc-matrix interfaces (c) In Fig 9, a scanning electron micrograph of a polished cross section normal to the gauge width direction is presented. The specimen is a tensile-fractured 90 nm-thick Pyc interphase com- posite. As can be seen, matrix cracking initiates preferentially from large interbundle pores resulting from overlay of the cross ing fiber bundles. Matrix cracks across the longitudinal fiber bundles were rarely observed at other locations. IV. Discussion The influence of Pvc on the mechanical properties of SiC/Sic composites were first studied for the CG-Nicalon"/PyC/CVI SiC system. For this system, Lowden reported that the peak flexural strength was achieved at a pvc of 200 nm, while the nterfacial shear stress measured by a micro-indentation method was inversely proportional to PVc. He concluded that the pri- 40m mary factor that controlled the macroscopic strength at ambient Fig & Longitudinal fiber bund\&/CVI-SiC composites. Fiber frac- fracture surfaces of Tyr. temperature was a mitigation of thermal residual stress, arising from coefficient of thermal expansion(CTE) mismatch between ure surfaces are shown in the inset. [ 0/90]. [0/90]: Pyc, pyrolytic the fiber and the matrix, and the resulting fiber clamping be- carbon; CVI, chemically vapor infiltrated cause of the presence of compliant PyC interlayer. He also noted that the Pyc coating protected the fiber from extensive chemical bonding with the matrix and chemical damage during cv to possess CtE very similar to that of CVI-Sic matrices cessing. Later, Singh et al.." correlated the flexural strength Tyranno-SA fiber consists primarily of p-SiC crystal grains the in situ fiber strength estimated from fracture mirror with a small amount(- 10% volume)of glassy carbon, which is analysis, for thin Pyc conditions. Very similar trends in CTEs of CG-Nicalon ain junctions in the core region. The and Hi-Nicalon" were 3.2 and system by Yang er ot ported for the Hi-Nicalon"/PyC/CVI-SiC tional stresses were 3.5×10-6/Kat25°-500°C, respectiv in contrast to the av. erage cte of~39×10-6/Kat25°-500 for CVDβsiC The characteristics of the Tyranno-SA/PyC/CVI-SiC sys- Second, it is unlikely that Tyranno -SA fibers experience sub- tem investigated in this study are significantly different from the stantial chemical damage during processing. Anticipated reac above cases. Most importantly, Tyranno-SA fiber is expected tions involving excess carbon and silica in CG-Nicalon" and Hi-

was noticed for the pull-out length among composites with a tPyC of 25–250 nm, as compared in Figs. 8(A) and (B). The composite with a tPyC of 830 nm exhibited substantially longer fiber pull-out lengths of 50–200 mm as shown in Fig. 8(C). De￾bonding of the longitudinal fiber interfaces occurred mostly at the interface between the PyC interphase and the matrix. The fiber fracture surfaces were dominated by grain boundary frac￾ture accompanied by river patterns, which is typical for Tyr￾annot-SA fibers. Most of the fiber fractures originated at the surface (insets to Fig. 8) but some originated within the core region. Side surfaces of the pull-out fibers appeared free from scratches or any evidence of wear damage, exhibiting only the original roughness because of the surface grain structures. As for the fracture surfaces of transverse fiber bundles, major cracking took place mostly within the fiber bundles, and the debond was observed at both the fiber–PyC and the PyC–matrix interfaces. In Fig. 9, a scanning electron micrograph of a polished cross section normal to the gauge width direction is presented. The specimen is a tensile-fractured 90 nm-thick PyC interphase com￾posite. As can be seen, matrix cracking initiates preferentially from large interbundle pores resulting from overlay of the cross￾ing fiber bundles. Matrix cracks across the longitudinal fiber bundles were rarely observed at other locations. IV. Discussion The influence of tPyC on the mechanical properties of SiC/SiC composites were first studied for the CG-Nicalont/PyC/CVI￾SiC system. For this system, Lowden reported that the peak flexural strength was achieved at a tPyC of B200 nm, while the interfacial shear stress measured by a micro-indentation method was inversely proportional to tPyC. 23 He concluded that the pri￾mary factor that controlled the macroscopic strength at ambient temperature was a mitigation of thermal residual stress, arising from coefficient of thermal expansion (CTE) mismatch between the fiber and the matrix, and the resulting fiber clamping, be￾cause of the presence of compliant PyC interlayer. He also noted that the PyC coating protected the fiber from extensive chemical bonding with the matrix and chemical damage during CVI processing. Later, Singh et al.,24 correlated the flexural strength with the in situ fiber strength estimated from fracture mirror analysis, for thin PyC conditions. Very similar trends in tPyC dependence of flexural strength and interfacial debond and fric￾tional stresses were reported for the Hi-Nicalont/PyC/CVI-SiC system by Yang et al. 25,26 The characteristics of the Tyrannot-SA/PyC/CVI-SiC sys￾tem investigated in this study are significantly different from the above cases. Most importantly, Tyrannot-SA fiber is expected to possess CTE very similar to that of CVI-SiC matrices. Tyrannot-SA fiber consists primarily of b-SiC crystal grains with a small amount (B10% volume) of glassy carbon, which is concentrated at the multigrain junctions in the core region. The CTEs of CG-Nicalont and Hi-Nicalont were 3.2 and 3.5
106 /K at 251–5001C, respectively,27 in contrast to the av￾erage CTE of B3.9
106 /K at 251–5001C for CVD b-SiC. Second, it is unlikely that Tyrannot-SA fibers experience sub￾stantial chemical damage during processing. Anticipated reac￾tions involving excess carbon and silica in CG-Nicalont and Hi￾Fig. 8. Longitudinal fiber bundle regions in fracture surfaces of Tyr￾annot-SA 2D Plain-Weave [0/90]/PyC/CVI-SiC composites. Fiber frac￾ture surfaces are shown in the inset. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Fig. 7. Influence of PyC interphase thickness on the initial tangential modulus of Tyrannot-SA 2D Plain-Weave [01/901]/(SiC/) PyC/ CVI￾SiC composites. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. 3092 Journal of the American Ceramic Society—Katoh et al. Vol. 88, No. 11

ovember 2005 Mechanical Properties of Thin Py C Interphase Sic-Matrix Composites 3093 fraction of a low modulus in correlated with tpyc. If we composite modulus. in wh of the interpha mal to, and 2 /3 along, the loading direction. it gitudinal fibers EEc E=31E+(-HE+(1-H)E(1) where Vi is the interphase volume fraction, and Ei and ee are the transverse cracks nterphase elastic modulus and the effective modulus of the combined structure of the fibers and matrix, respectively. The model assumes that one-half of the transversal fiber interphase is in parallel to the loading direction and the other half is perpen- dicular to that. This assumption does not strictly represent the d inter bundle pore case of an axisymmetrical interphase but should be sufficient for a rough estimation of the interphase stiffness. The fitting of eq elds Ei=x 10 GP of a polished cross section normal to the GPa for the ORNL Tyranno-SA/PyC/CVI-SiC system. The 090. [0/90]: PyC, pyrolytic carbon; CVI, 18 GPa for isotropic PyC.On the other hand, CTE of chemically vapor tropic PyC is reportedly -5x 10-6 K- at 0% C.30Be- cause of the very small elastic modulus, very thin PyC layer, and small CTE mismatch against B-SiC, it is reasonable to ignore the Nicalon fibers and the reactants (MTS and hydrogen) do not effect of fiber clamping by the pyc interphase ditionally, the mismatch in elastic modulus between the fibers As the cte mismatch effect on fiber clamping is probably and matrix is small. as the manufacturer-claimed young,s mod- negligible, the fiber shear debond stress is determined by the ulus of Tyranno-SA is 400 GPa in contrast to 220 GPa fo chemical bonding and static friction at the interface. Only the The interesting features shown in Fig. 5 are the general in- primary factor that controls the frictional stress is the interphase sensitivity of UTS to tPyc at tpyc300 debond of interfaces near the surface during sample preparation thick PyC interphase is not of major concern. As shown in Fig. 7, the tangential elastic modulus decreases For further investigation of interfacial properties, the method slightly as tpyc increases. This is an effect of the larger volume of unloading/reloading hysteresis analysis proposed by Agag- 8 A NIMs Tyranne-SA ◆ Ty ranno-sA/PYCNCVi-sc 293K A NIMS HHNiealon/PyCICVI-SIC Berkovich (Yang) a 80o DoRNL CG-Nicalon/P c/cvi-sic. Vickers (Lowden) 0.2 ◆ 293K 00 04 Reciprocal PyC Interphase Thickness(nmr) Fig. ll. Interfacial fiber debond load divided by the fiber surface area Fig. 10. Influence of PyC interphase thickness on strain to fracture of for Tyranno-SA 2D Plain-Weave [o/90/PyC/CVI-SiC composites as yranno-SA 2D Plain-Weave [o/90(SiC/ PyC/cVI-SiC composites. measured by the single-fiber push-out method. [090). [0/90]: PyC. 0/90.10/90: PyC, pyrolytic carbon; CVI, chemically vapor infiltrated rolytic carbon; CVI, chemically vapor infiltrated

Nicalont fibers and the reactants (MTS and hydrogen) do not occur on the stoichiometric SiC surface of Tyrannot-SA.28 Ad￾ditionally, the mismatch in elastic modulus between the fibers and matrix is small, as the manufacturer-claimed Young’s mod￾ulus of Tyrannot-SA is 400 GPa in contrast to 220 GPa for CG-Nicalont. The interesting features shown in Fig. 5 are the general in￾sensitivity of UTS to tPyC at tPyCt300 nm and the significantly higher UTS at tPyC\600 nm. The former can be attributed pri￾marily to the relative unimportance of the role of interphase in mitigating the fiber-clamping stress arising from the CTE mis￾match. The latter observation, along with the longer fiber pull￾out length and the greater strain to fracture at the thicker in￾terphase shown in Figs. 8 and 10, indicates that an interphase thicker than B300 nm allows interfacial sliding necessary to obtain the highest UTS. The primary reason for the thicker op￾timum tPyC for the highest strength in the Tyrannot-SA com￾posites compared with that in the CG-Nicalont and Hi￾Nicalont composites23,25 is presumably because of the differ￾ence in fiber surface roughness. It is not possible to further nar￾row down the optimum tPyC range for the highest UTS from the present data. However, an optimum tPyC may be found at 4300 nm for applications in which the environmental performance of thick PyC interphase is not of major concern. As shown in Fig. 7, the tangential elastic modulus decreases slightly as tPyC increases. This is an effect of the larger volume fraction of a low modulus interphase, because the porosity is not correlated with tPyC. If we assume a very simple model for the composite modulus, in which 1/3 of the interphase layer is nor￾mal to, and 2/3 along, the loading direction, it would be: Ec ¼ 1 3 EiEe ViEe þ ð1 ViÞEi þ 2 3 ½ ViEi þ ð1 ViÞEe (1) where Vi is the interphase volume fraction, and Ei and Ee are the interphase elastic modulus and the effective modulus of the combined structure of the fibers and matrix, respectively. The model assumes that one-half of the transversal fiber interphase is in parallel to the loading direction and the other half is perpen￾dicular to that. This assumption does not strictly represent the case of an axisymmetrical interphase but should be sufficient for a rough estimation of the interphase stiffness. The fitting of Eq. (1) to the data in Fig. 7 yields Ei 5B10 GPa and Ee 5B250 GPa for the ORNL Tyrannot-SA/PyC/CVI-SiC system. The estimated Ei value is slightly lower but close to the reported 12– 18 GPa for isotropic PyC.29 On the other hand, CTE of iso￾tropic PyC is reportedly B5
106 K1 at 01–10001C.30 Be￾cause of the very small elastic modulus, very thin PyC layer, and small CTE mismatch against b-SiC, it is reasonable to ignore the effect of fiber clamping by the PyC interphase. As the CTE mismatch effect on fiber clamping is probably negligible, the fiber shear debond stress is determined by the chemical bonding and static friction at the interface. Only the frictional stress should be affected by the variation in tPyC. The primary factor that controls the frictional stress is the interphase compliance. Yang et al. 22 evaluated the effect of tPyC on the fiber debond initiation load during Berkovich indentation in Tyran￾not-SA/PyC/CVI-SiC composites (produced at NIMS, Japan), Q4 which are similar to the composites studied in this work. Their result, replotted in Fig. 11, showed that the debond load was nearly independent of tPyC in the range of 50–100 nm. As tPyC increased beyond B100 nm, the debond load gradually de￾creased through mitigation of the locking effect arising from the fiber surface roughness. Similar plots have been placed together in Fig. 11 for NIMS Hi-Nicalont/PyC/CVI-SiC26 and ORNL CG-Nicalont/PyC/CVI-SiC composites.28 The tPyC effect is ap￾parently similar to that in the Tyrannot-SA composites at \100 nm in spite of a likely different cause of the friction. The measured anomalously low debond stresses for the Hi-Ni￾calont composites at tPyCt100 nm are likely a result of partial debond of interfaces near the surface during sample preparation where large thermal residual shear stress existed. For further investigation of interfacial properties, the method of unloading/reloading hysteresis analysis proposed by Vagag￾Fig. 9. Optical micrograph of a polished cross section normal to the gauge width direction in Tyrannot-SA 2D Plain-Weave [0/90]/PyC (90 nm)/CVI-SiC composites. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Fig. 10. Influence of PyC interphase thickness on strain to fracture of Tyrannot-SA 2D Plain-Weave [0/90]/(SiC/) PyC/CVI-SiC composites. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Fig. 11. Interfacial fiber debond load divided by the fiber surface area for Tyrannot-SA 2D Plain-Weave [0/90]/PyC/CVI-SiC composites as measured by the single-fiber push-out method. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Q1 November 2005 Mechanical Properties of Thin PyC Interphase SiC–Matrix Composites 3093

Journal of the American Ceramic Society--Katoh et Vol. 88. No. I IL Tyranno-SA/ PyC160nm)/CVl-SC ◆ Tanno- SA/PyCKCV-sc 293K 是 2=100MPa PyC Interphase Thickness(nm) Reloading Stress(MPa) Fig 14. Thermal misfit stress plotted against PyC interphase thickness Fig 12. Reloading reciprocal moduli measured during tensile testing for Tyranno-SA 2D Plain-Weave [0/90V(SiCD PyC/CVI-SiC compos- posite. [0/90]. [0%90%] PyC Duve, lo/90/A-r sequences. The material is ites. [0/90]. [0/900]: PyC, pyrolytic carbon; CVI, chemically vapor infil- 2D C(160 nm)/CVI-SiC com- ytic carbon; CVI, chemically vapor infiltrated nearly constant for 60 nm<tpyC<300 nm, while it was signif gini and co-workers was applied. 31,32 All the tested composites icantly higher at $60 nm and lower at 2 600 nm. As below exhibited hysteresis curves, which are characteristic of materials 60 nm the surface roughness of fiber presumably dominate with a relatively large debond energy. The reloading load-strain the sliding stress, Ipyc greater than this would be suggested in curve comprises of the initial parabolic segment that represents order to achieve optimal composite properties. The threshold dynamic frictional stress(sliding stress) followed by a linear Ipyc might be affected by the Sic overcoating onto the fiber, segment, as shown in Fig. 12, where reciprocal moduli during although this could not be confirmed from the present data. The reloading are plotted against apparent stress after unloading threshold at 60 nm is in contrast to the observation that the from different peak stresses (o,). As nearly constant sliding initial debond stress as measured by push-out was constant for stress parameters were obtained when op2 150 MPa, the fol- c<100 nm and decreased at higher pyc values. The t' plot lowing discussion is based on analysis of reloading behavior af- Fig. 13 primarily shows the influence of compliant layer thick- ter unloading from x 200 MPa ess on the dynamic friction of rigid and uncorrelated rough According to the model by vagaggini et al, the initial slope in surfaces on both its sides, while the push-out stress is related to ig. 12 gives the inelastic strain inde debond and slide initiation of completely mated rough surfaces Assuming an average matrix crack spacing of 0.75 mm, which corresponds to one-half of the tow interval, the dynamic fric- tional stresses estimated from Fig. 13 are approximately 3-10 MPa. the low end of the estimated frictional stresses for the where a; and bi are the Hutchinson-Jensen parameters, f is the present composites is similar to the interfacial sliding stress as olume fraction in the loading direction, t is the interfacial measured with a fat-bottomed indenter for CG-Nicalon"/Pyc/ stress. R is the fiber radius. and d is the average matrix CVI-SiC composites with interphases sufficiently thick to miti- spacing. Em is the effective modulus of the entire com- gate the clamping stress. This suggests that the frictional stress posite excluding the longitudinal fibers and calculated from approaches similar values when a thick Pyc interphase is ap- mposite modulus Ec, fiber modulus Er, and f. Eq(2), plied regardless of the primary cause of friction without the the sliding stress parameter t'=td/R is plotted tion of thick interphase tpyc in Fig. 13. Assuming similar matrix crack for all Thermal residual stress was also estimated by following the composites at the same stress level, the sliding stress remained method proposed by vagaggini et al. The result is presented misfit stress oas a function of Pvc in Fig. 14. The fell in the range of 10-20 MPa for most specimens, regardless of tpyc. The 293K axial stress in matrix is in tension and roughly comparable with ◆ Tyranno-sA/PyC/CVI-SIC o' in magnitude for this material system. The conversion misfit stress to residual stress components is provided else- where. The existence of a small tensile residual stress in the matrix is supported by the fact that the extrapolated linear por tions of the regression lines intersect at -30 to -10 MPa in the stress-strain charts. These numbers are about an order of mag- nitude lower than that reported for a CG-Nicalon"/PyC/CV SiC composite. Along with the observed insensitivity of ele vated temperature tensile strength to Pvc, a very small CtE mismatch stress for the present composite system is confirmed. The influence of PVc in the range of Pyc Interphase Thickness(nm) properties of the Tyranno-SA /PyC/CVI-SiC composite system was generally very minor. Thin tpyc such as x50 nm did not Fig 13. Influence of pyrolytic carbon(PyC) interphase thickness or result in a significant deteriorating effect on composite proper dynamic interfacial sliding stress parame ties, which has commonly been reported for conventional SiC

gini and co-workers was applied.31,32 All the tested composites exhibited hysteresis curves, which are characteristic of materials with a relatively large debond energy. The reloading load–strain curve comprises of the initial parabolic segment that represents dynamic frictional stress (sliding stress) followed by a linear segment, as shown in Fig. 12, where reciprocal moduli during reloading are plotted against apparent stress after unloading from different peak stresses (sp). As nearly constant sliding stress parameters were obtained when sp\150 MPa, the fol￾lowing discussion is based on analysis of reloading behavior af￾ter unloading from B200 MPa. According to the model by Vagaggini et al., the initial slope in Fig. 12 gives the inelastic strain index l ¼ b2ð1 a1fÞ 2 4f 2tEm R d (2) where ai and bi are the Hutchinson–Jensen parameters,33 f is the fiber volume fraction in the loading direction, t is the interfacial sliding stress, R is the fiber radius, and
d is the average matrix crack spacing. Em is the effective modulus of the entire com￾posite excluding the longitudinal fibers and calculated from composite modulus Ec, fiber modulus Ef, and f. Using Eq. (2), the sliding stress parameter t0  t
d=R is plotted as a function of tPyC in Fig. 13. Assuming similar matrix crack spacing for all composites at the same stress level, the sliding stress remained nearly constant for 60 nmotPyCo300 nm, while it was signif￾icantly higher at t60 nm and lower at \600 nm. As below B60 nm the surface roughness of fiber presumably dominates the sliding stress, tPyC greater than this would be suggested in order to achieve optimal composite properties. The threshold tPyC might be affected by the SiC overcoating onto the fiber, although this could not be confirmed from the present data. The threshold at B60 nm is in contrast to the observation that the initial debond stress as measured by push-out was constant for tPyCo100 nm and decreased at higher tPyC values. The t0 plot in Fig. 13 primarily shows the influence of compliant layer thick￾ness on the dynamic friction of rigid and uncorrelated rough surfaces on both its sides, while the push-out stress is related to debond and slide initiation of completely mated rough surfaces. Assuming an average matrix crack spacing of 0.75 mm, which corresponds to one-half of the tow interval, the dynamic fric￾tional stresses estimated from Fig. 13 are approximately 3–10 MPa. The low end of the estimated frictional stresses for the present composites is similar to the interfacial sliding stress as measured with a flat-bottomed indenter for CG-Nicalont/PyC/ CVI-SiC composites with interphases sufficiently thick to miti￾gate the clamping stress.34 This suggests that the frictional stress approaches similar values when a thick PyC interphase is ap￾plied regardless of the primary cause of friction without the thick interphase. Thermal residual stress was also estimated by following the method proposed by Vagaggini et al. 31 The result is presented in misfit stress sT as a function of tPyC in Fig. 14. The sT fell in the range of 10–20 MPa for most specimens, regardless of tPyC. The axial stress in matrix is in tension and roughly comparable with sT in magnitude for this material system. The conversion of misfit stress to residual stress components is provided else￾where.31 The existence of a small tensile residual stress in the matrix is supported by the fact that the extrapolated linear por￾tions of the regression lines intersect at 30 to 10 MPa in the stress–strain charts. These numbers are about an order of mag￾nitude lower than that reported for a CG-Nicalont/PyC/CVI￾SiC composite.32 Along with the observed insensitivity of ele￾vated temperature tensile strength to tPyC, a very small CTE mismatch stress for the present composite system is confirmed. V. Conclusions The influence of tPyC in the range of B50–B300 nm on tensile properties of the Tyrannot-SA/PyC/CVI-SiC composite system was generally very minor. Thin tPyC such as B50 nm did not result in a significant deteriorating effect on composite proper￾ties, which has commonly been reported for conventional SiC Fig. 12. Reloading reciprocal moduli measured during tensile testing incorporating multiple unloading–reloading sequences. The material is a Tyrannot-SA 2D Plain-Weave [0/90]/PyC (160 nm)/CVI-SiC com￾posite. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infiltrated. Fig. 13. Influence of pyrolytic carbon (PyC) interphase thickness on the dynamic interfacial sliding stress parameter. Fig. 14. Thermal misfit stress plotted against PyC interphase thickness for Tyrannot-SA 2D Plain-Weave [0/90]/(SiC/) PyC/CVI-SiC compos￾ites. [0/90], [01/901]; PyC, pyrolytic carbon; CVI, chemically vapor infil￾trated. 3094 Journal of the American Ceramic Society—Katoh et al. Vol. 88, No. 11

ovember 2005 Mechanical Properties of Thin PyC Interphase SiC-Matrix Composites fiber-CVI-SiC composites. The insensitivity of the composites properties to the interphase thickness will be of great advantage Silicon Carbide. ". Nucl. Mater. 307-311 1141-5(2002) 429 in Carbon For tpyc 600 nm at the expense of composite rigidity erties in FCVI SIC Ma oichiometric SiC Fiber Composite System, J.Macl. Mater,307-311,1205-90002). cations in which a thick py c interphase is allowed and rigidity is Efets i esi p ratti r sicsic a an comarae-duation s sp ir not of major concern. In conclusion, the Tyranno-SA/PyC/CVI-Sic composites for Fusion Energy Application, April 12-13. 2002. International Energy Agency, with thin(25 nm) PyC interphases appear to be good candi San Diego, CA, pp 74-86 dates for nuclear applications. Pyc interphases thicker than 60 nm will be preferred for optimal composite property. The 305 in Small Specimens Test Techiniques: Fourth Volume, ASTA STP 1418. Edited upper limit of interphase thickness will depend on the condition by M. A Sokolov,J. D. Landes, and G. E. Lucas. ASTM International, West 2 K Hironaka T Nozawa, T Hinoki N. Igawa, Y Katoh. L L. Sneadand A Kohyama, "" High-Temperature Tensile Strength of Near-Stoichiometric SiC/SiC Acknowledgments Y. Yan, C. W. Chen. P. C. Fang. K. M. Yin, F.R. Chen. Y. Katoh. A. The authors would like to thank Drs. Edgar Lara-Curzio. Hua-Tay Lin R chard A Lowden and Steven J. Zinkle at ORNL. Dr Wen Y ang at NIMS. D SicSiC Composites with ESI Method, "J. Nucl. Mater, 329-333, 513-7(200 Kyoto University, and the revi ang, A. Kohyama, Y. Katoh, H. Araki, J. useful advice and discussion. The CVI processing was performed by Drs. Tom- Carbon and Silicon Carbide/Carbon Interlayers on the Mechanical Behavior itsugu Taguchi and Naoki Igawa at JAERI and Mr. Jerry McLaughlin at ORNL. of Tyranno-SA-Fiber-Reinforced Silicon Carbide-Matrix Composites,"J.Anm. Ceran.Soe,86.851-6(2003) 2R. A. Lowden, "Fiber Coatings and the Mechanical Properties of a Fibe einforced Ceramic 2D. Singh, J. P Singh, and M. J. Wheeler, "Mechanical Behavior of SiC(D/SiC R. Naslain, "Design, Preparation and Properties of Non-Oxide CMCs for Ap- ation in Engines and Nuclear Reactors: An Overview, Compos. Sci. Technol. 2w. Yang H. Araki. T Noda. J Y. Park, Y Katoh. T Hinoki. J. Yu andA w. R. Corwin, L.L. Snead. S.J. Zinkle, R. K. Nanstad. A. F. Rowcliffe L.K. Kohyama.""Hi-Nicalon Fiber-Reinforced CVI-SiC Matrix Composites: I Ef- sur, R. W. Swindeman, W. Ren, D. F. wilson, T. E. McGreevy, P. L. ects of Py C and PyC-SiC Multilayers on the Fracture Behaviors and Flexural and R&D Needs to Assess Viability", Oak Yang, H. Araki, A. Kohyama. Y. Katoh, Q. Hu, H. Suzuki, and T Noda, Laboratory Technical Memorandum, U.S. Department of Energy Hi-Nicalon" Fiber-Reinforced CVI-SiC Matrix Cor Flexural Properties n,43.257477 BRam上和mA购 Y Katoh, A KohyamaR且Jmg(、、 rind Control of the Fiber-Matrix Interface in Fusion Reactors,J. Nucl. Mater. 329-333. 56-65(2004) 0431526,pp.105 ng. D. K. Sze. M. billo Lowden“Cha toslavsky, and ARIES Team, " Fusion Power Core Engineering for the Ceramic Matrix Composites": ORNL/TM-1103, Oak Ridge National Laborato- D. T Ingersoll, C. w. Forsberg, L J Williams. J. P. Renier. D. F. 2R. J. Price and J. C. Bokros. "Mechanical Properties of Neutron-Irradiated Wilson, S.J. Ball, L. Reid. w.R. Corwin G. D. Del Cul, P. F. Peterson, H. Zha Pyrolytic Carbon, "J. Nucl. Mater 21, 158( 1967) P.S. Pickard, E J. Parma, and M. Vernon. ""Status of Preconceptual Design of 30F. Ho."Material models Pyrocarbon and Pyrolytic Silicon Carbide he Advanced High Temperature Reactor(AHTR) ORNLTM-2004/104: Oak CEGA-002820, CEGA Corporation, 3E. Vagaggini, J-M. Domergue, and A. G. Evans, "Relationships Between R.J. Price, "" Properties of Silicon Carbide for Nuclear Fuel Particle Coatings, Ceramic matrix co Nucl. Mater. 219. 70-86 L. L. Snead. Y. Katoh. A. Kohyama, J. L. Bailey, N.L. Vaughn, andR. A 33J. w. Hutchinson and H. M. Jensen, ""Models of Fiber-Debonding and Pull- Lowden,“ Evaluation ometric Silicon Carbide The Effect of fiber Coating Thickness on the Interfacial Properties of a Continuous Fiber-Reinforced Ceramic Composites for Fusion Applications. Fusion Sci. Technol. 44. 155-62 eramic Composite, Ceran. Eng. Sci. Proc. 15. 989-1000(1994 of-Tested and Neu- tron-Irradiated Silicon Carbide, J. Nucl. Mater, 108& 109, 732-8(1982)

fiber–CVI-SiC composites. The insensitivity of the composites’ properties to the interphase thickness will be of great advantage for manufacturing. For tPyCt50 nm, the tensile strength of the composites may possibly be slightly lowered, accompanied by a slight decrease in strain to fracture. A substantial increase in dynamic interfacial sliding stress at tPyCt60 nm was noted. Therefore, an excess interfacial friction may have caused the potential strength re￾duction in this case. For tPyC\60 nm, the primary tPyC effects were reduced composite rigidity and increased interlayer com￾pliance. They caused a slightly negative tPyC dependence of pro￾portional limit stress, tangential modulus, and dynamic sliding stress, as well as a strong tPyC dependence of the static frictional stress upon the initial debonds. The strain to failure exhibited the reverse trend of the tangential modulus. A significant in￾crease in the ultimate tensile strength and fracture to strain was observed for tPyC4600 nm at the expense of composite rigidity. An optimum tPyC may be found in this range for certain appli￾cations in which a thick PyC interphase is allowed and rigidity is not of major concern. In conclusion, the Tyranno-SA/PyC/CVI-SiC composites with thin (\25 nm) PyC interphases appear to be good candi￾dates for nuclear applications. PyC interphases thicker than B60 nm will be preferred for optimal composite property. The upper limit of interphase thickness will depend on the condition of neutron irradiation. Acknowledgments The authors would like to thank Drs. Edgar Lara-Curzio, Hua-Tay Lin, Ri￾chard A. Lowden, and Steven J. Zinkle at ORNL, Dr. Wen Yang at NIMS, Drs. Tatsuya Hinoki and Akira Kohyama at Kyoto University, and the reviewers for useful advice and discussion. The CVI processing was performed by Drs. 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