ournal Am Cerum,soc,879120-17252004 Process and Mechanical Properties of in Situ Silicon Carbide-Nanowire-Reinforced Chemical Vapor Infiltrated Silicon Carbide/Silicon Carbide Composite Wen Yang, Hiroshi Araki, Akira Kohyama, *. Somsri Thaveethavorn, Hiroshi Suzuki, and Tetsuji Noda National Institute for Materials Science, Tsukuba 305-0047, Japan Institute of Advanced Energy. Kyoto University, Uji, Kyoto 611-0011, Japan A SiC nanowire/Tyranno-SA fiber-reinforced SiC/SiC com- transverse matrix cracks propagate easily through those fiber posite was fabricated via simple in situ growth of SiC nano- bundles perpendicular to the direction of the main stress in the wires directly in the fibrous preform before CVI matrix materials on either tensile or flexural external force, especially densification; the purpose of the SiC nanowires was to mark when these bundles are insufficiently deposited within the matrix edly improve strength and toughness. The nanowires consisted ak points tes and mak of single-crystal B-phase Sic with a uniform -5 nm carbon little contribution to strength and toughness. Challenges for further shell; the nanowires had diameters of several tens to one improvements of strength and toughness of SiC/SiC composites hundred nanometers. The volume fraction of the nanowires in should follow on strengthening and toughening these weak points the fabricated composite was -5%. However, the composite as well as the brittle matrix did not show significant increase in strength and toughness One-dimensional nanostructures, such as SiC nanowires and likely because of strong bonding between the nanowires and carbon nanotubes, have great potential for use in composite the matrix caused by the very thin carbon coating on the nano- materials as reinforcements because of their significantly greater wires. Little debonding and pullout of SiC nanowires from the strength than their bulk counterparts. 0- Several efforts have matrix were observed at the fracture surfaces of the composite. been made to fabricate carbon-nanotube-reinforced compos ites. 3. However, a significant reinforcement effect of the strong and expensive nanotubes on the mechanical properties has not been observed, likely because of several critical difficulties, as HERE has been strong interest in ceramic-matrix composites for roposed by Zhan et al. First, the properties of the individual nany decades for a variety of high-temperature, high-stress carbon nanotube must be optimized. Second, the carbon nanotubes applications in aerospace, hot engine, and energy conversion must be sufficiently bonded to the matrix so that they actuall devices, The low fracture toughness of the ceramics can be carry the load. Third, the load must be distributed throughout the readily improved by the incorporation of reinforcement materials, nanotubes to ensure that the outermost layer does not shear off such as continuous -SiC-fiber-reinforced SiC-matrix (SiC/SiC) SiC nanowires also have been suggested as good reinforcement opposites . 6 In a SiC/SiC composite, a transverse matrix crack materials for ceramic-matrix composites. o The reported elastic can be deflected with energy dissipation that occurs via debonding modulus and ultimate bending strengths of SiC nanorods are at the fiber/matrix interface, crack deflection, crack bridging by the 610-660 GPa and 53.4 GPa, respectively. These values are much fibers, fiber sliding, and eventual fiber fracture. These energy. larger than those of SiC-based fibers, including the advanced dissipating mechanisms provide for improved apparent fracture Tyranno-SA fibers, which are newly developed Sic fibers and toughness and result in a noncatastrophic mode of failure. Many widely used for advanced Sic/SiC composites. Although several Sic/SiC composites for the purpose of improving strength and wires. no effort has been reported in the current literature on fracture tolerance, and a good understanding of the results has been Sic-nanowire-reinforced composites. This is possibly because of established the substantial challenges on homogeneity that disperse the Although the fracture tolerance of bulk Sic can be readily nanowires in the matrix with sound nanowire/matrix interfacial gains in fracture tolerance are basically from the fiber/matrix matrix, such as SiC, is likely to produce strong interfacial bonding, interfacial debonding and the bridging/deflection of transverse which results in unnecessarily improved strength over the pure matrix cracks by the fibers. The SiC matrix displays a brittle matrix. Hence, a compliant coating on the nanowires is necessary behavior similar to its bulk counterparts. In addition, when the for modified bonding between the nanowires and the matrix. composite is reinforced with two- or three-dimensional fabrics, Chemical vapor deposition( CVD) is a suitable and widely used method to produce SiC of various shapes of thin films, powders whiskers and nanorods 8-20 Kirchner and Knoll have obtained microscale CVD-SiC whiskers through thermal decomposition of T M Besmann--contributing editor CH, SiCI, (MTS), which is carried by hydrogen. Lespiaux et al. have performed a detailed study on the correlation between ga phase and supersaturation, nucleation process, and physicochem cal characteristics of CVD-SiC deposited from the MTS-H2 system; they have found that the supersaturation, Y and the type of Culture, Sports, Science, and Technology, based on screening and counseling by the kinetic process are determining factors for the control of the Atomic Energy Commission. morphology and microstructure of the deposited SiC. Based on this rican Ceramic Society National Institute for Materials Science knowledge, a simple CVD process has been developed, and B-SiC "Institute of Advanced Energy nanowires that have a thickness of -100 nm and length of several
journal / Im Ceram. Soc.. 87 [91 1720-17:5 12004) Process and Mechanical Properties of in Situ Silicon Carbide-Nanowire-Reinforced Chemical Vapor Infiltrated Silicon Carbide/Silicon Carbide Composite Wen Yang/ Hiroshi Araki,' Akira Kohyama.** Somsri Thaveethavom.^ Hiroshi Suzuki/ and Tetsuji. Noda' National Institute for Materials Science, Tsukuba 305-0047, Japan Institute of Advanced Energy. Kyoto University, Uji, Kyoto 611-0011, Japan A SiC nanowire/Tyranno-SA fiber-reinforced SiC/SiC composite was fabricated via simple in situ growth of SiC nanowires directly in the fibrous preform hefore CVI matrix densification; the purpose of the SiC nanowires was to markedly improve strength and toughness. The nanowires consisted of single-crystal p-phase SiC with a uniform —5 nm carbon shell; the nannwires had diameters of several tens to one hundred nanometers. The volume fraction of the nanowires in the fabricated composite was —5%. However, the composite did not show significant increase in strength and toughness, likely because of strong bonding between the nanowires and the matrix caused by the very thin carbon coating on the nanowires. Little debonding and pullout of SiC nanowires from the matrix were observed at the fracture surfaces of the composite. I. Introduction T HF.RE has been strong interest in ceramic-matrix composites for many decades for a variety of high-temperature, high-stress applications in aerospace, hot engine, and energy conversion devices.'""" The low fracture toughness of the ceramics can be readily improved by the incorporation of reinforcement materials, such as continuous-SiC-t1ber-reinforced SiC-matrix (SiC/SiC) composites.-^'' In a SiC/SiC composite, a transverse matrix crack can be deflected with energy dissipation that occurs via debonding at the fiber/matrix interface, crack deflection, crack bridging by the fibers, fiber sliding, and eventual flber fracture,'"^ These energydissipating mechanisms provide for itnproved appareni fracture toughness and result in a noncatastrophic mode of failure. Many efforts have been made to upgrade the interfacia! properties of SiC/SiC composites for the purpose of improving strength and fracture tolerance, and a good understanding of the results has been established.''"'' Although the fracture tolerance of bulk SiC can be readily improved by the incorporation of SiC reinforcement flbers, such gains in fracture tolerance are basically from the flber/matrix interfacial debonding and the bridging/deflection of transverse matrix cracks by the fibers. The SiC matrix displays a brittle behavior similar to its bulk counterparts. In addition, when the composite is reinforced with two- or three-dimensional fabrics. T. M. Besniann—contributing editor ManusL-ripi No. 10681. Received November 26. 2003; approved May 25. 2CK)4. Supported by the Budget lor Nuclear Research of the Ministry of Education, Culture, Sports, Science, and Technology, based on screenittg and counseling by the Atotiiic Energy Commission. 'Member. American Ceramic Society. ^National Institute for Materials Science. *lnstilute of Advanced Energy. transverse matrix cracks propagate easily through those flber bundles perpendicular to the direction of the main stress in the materials on either tensile or flexural external force, especially when these bundles are insufficiently deposited within the matrix. These fiber bundles are weak points in the composites and make little contribution to strength and toughness. Challenges for further improvements of strength and toughness of SiC/SiC composites should follow on strengthening and toughening these weak points as well as the brittle matrix. One dimensional nanostructures, such as SiC nanowires and carbon nanotubes, have great potential for use in composite materials as reinforcements because of their significantly greater strength than their bulk counterparts.'""'" Several efforts have been made to fabricate carbon-nanotubc-reinforced composites.'"^"'* However, a significant reinforcement effect of the strong and expensive nanotubes on the mechanical properties has not been observed, likely because of several critical difflculties, as proposed by Zhan et ai^^ First, the properties of the individual carbon nanoiube must be optimized. Second, the carbon nanotubes must be sufficiently bonded to the matrix so that they actually carry the load. Third, the load must be distributed throughout the nanotubes to ensure that the outermost layer does not shear off. SiC nanowires also have been suggested as good reinforcement materials for ceramic-matrix composites.'" The reported elastic modulus and ultimate bending strengths of SiC nanorods are 610-660 GPa and 53.4 GPa, respectively. These values are much larger than those of SiC-based fibers, including the advanced Tyranno-SA flbers. which are newly developed SiC fibers and widely used for advanced SiC/SiC composites.'^ Although several methods have been developed for fabrication of SiC nanowires,'^'"* no effort has been reported in the current literature on SiC-nanowire-reinforced composites. This is possibly because of the substantial challenges on homogeneity that disperse the nanowires in the matrix with sound nanowire/matrix interfacial bonding strength. Direct burying of the SiC nanowires in a ceramic matrix, such as SiC, is likely to produce strong interfacial bonding, which results in unnecessarily improved strength over the pure matrix. Hence, a compliant coating on the nanowires is necessary for modified bonding between the nanowires and the matrix. Chemical vapor deposition (CVD) is a suitable and widely used method to produce SiC of various shapes of Ihin films, powders, whiskers, and nanorods.'" ^" Kirchner and Knoll"' have obtained microscale CVD-SiC whiskers through thermal decomposition of CHiSiCli (MTS), which is carried by hydrogen. Lespiaux et a(}~ have performed a detailed study on the correlation between gas phase and supersaturation, nucleation process, and physicochemical characteristics of CVD-SiC deposited from the MTS-H^ system; they have found that the supersaturation, •y. and the type of kinetic process arc determining factors for the control of the morphology and microstructure of the deposited SiC. Based on this knowledge, a simple CVD process has been developed, and p-SiC nanowires that have a thickness of —100 nm and length of several 1720
September 2004 CVI Furnace Process and Mechanical Properties of in Situ SiC Nanowire-Reinforced CVI Sic/SiC Composite mber of the geometry. The stack was located in the vertical hot chan Hot Chamber CVI furnace, as shown in Fig. 1. Before the nanowire-growing process was begun. the furnace(preform) was heated to 1223 K under vacuum(<0. I Pa) using a rotary pump. The rotary pump Composite Preform then was replaced by a diaphragm pump, and CH, was introduced Carbon Mesh Gas Distributor nto the specimen chamber at a constant flow rate of 200 sccm standard cubic centimeters per minute) to deposit a carbon coating on the fibers in the preform as the fiber/matrix interlayer. The Valve exhausted gas was evacuated using the diaphragm pump. During H the CVD carbon-coating process, the pressure in the chamber was H2 carrier maintained at 14.7 kPa (by adjusting the evacuation rate of the Cold Diaphragm gm pump). After the specimen was processed for 70 min Trap Pump the CHa flow was stopped, and the system was maintained at 1223 K for 10 min to make certain of complete reaction of the residual CHa in the chamber. The specimen then was heated to 1373 K and Temperature b od MTS Bath Rotary MTS was introduced into the chamber to grow SiC nanowires in the preform. The MTS was carried by hydrogen-gas using a typical bubble system(Fig. I). A bypass hydrogen flow also was used to Fig. 1. Schematic diagram of the CvI system dilute the MrS concentration in the chamber. During this process. the MTS bath temperature was maintained at 293 K. The flow rates of carrier and bypass hydrogen-gas were 20 and 1000 sccm. tens of micrometers have been produced in our recent work 23 respectively. Therefore, the flow rate of the MTS was 1.8X 10 mol/min. The nanowire process was conducted for 2 h at a Chemical vapor infiltration (CVD) essentially has been developed decreased specimen chamber pressure of 1. 47 kPa. The MTS and preform held at an elevated temperature. Matrix materials are hydrogen-gas flows then were stopped. After the system wa maintained at this level for- 10 min, it was cooled to deposited on the substrate structure via a standard CVD reaction temperature, and the weight of the preform was measured to Because of such advantages, CVI is widely applied for the estimate the amount of SiC nanowires in the preform. One fabric fabrications of fiber/whisker reinforced/ceramic matrix compos- sheet from the center of the preform was taken out for microscopic ites, including SiC/SiC composites examination before the synthesized SiC nanowires in the preform The SiC nanowires in our previous work were a plain-woven Tyranno-SA fabric sheet using a typical C\ were the same as described before but for a shorter processing system. Here we report, for the first time, the fabrication of a Sic time, i.e., 20 min. This carbon coating was supposed to decrease growth of Sic nanowires in the composite preform before the this carbon-coating process, another fabric sheet was used to the nanowires to modify nanowire/matrix interfacial bonding. The After the completion of the second carbon-coating process, the main interests were the flexural strength and fracture toughness of preform was compressed with a set of graphite fixtures to 2.2 mm the composite and the effects of the SiC nanowire in thickness, which resulted in a volume fraction of fibers of -43%6. Finally, the preform was densified with SiC matrix IL. Experimental Procedures material for 18 h using a CVI process at 1273 K under 14. 7 kPa The total hydrogen flow was 1000 sccm, and the MTS: hydrogen I) Processes for Carbon-Coated SiC Nanowires and volume ratio was 1: 10. The composite was named NF-C SiC/SiC Composites (nanowire- and fiber-reinforced composite) The nanowires were synthesized directly in a compo A conventional Tyranno-SA/SiC composite, named F-C. also form using a CVI system, as schematically shown in Fig was fabricated for comparison. A similar CVI process as described composite preform was prepared by stacking 14 sheets (40 above was used, except the nanowire portion of the process was diameter) of plain-woven Tyranno-SA fiber cloths in a deleted a) 20m 5u m Fig. 2. SiC nanowires on the Tyranno-SA fabric sheet
September 2004 Process and Mi'cbunical Properties oj in Sini SiC Nanowire-Reinforced CVl SiC/SiC Camposile 1721 C\'l Furnace Hot Chamber Tlienitocouple Composite Preform Carbon Mesh Gas Distributor Klowmeler, Valve . CH4 H2 by pass H2 carrier Temperature Controller fiii? t t t Diaphragm Pump MTS Bath Rotary Pump Fig. 1. Schematic diagram of the CVI system. tens of micrometers have been produced in our recent Chemical vapor infiltration (CVI) essentially has been developed from CVD.""^'^ In CVI. gaseous reactants infiltrate a porous preform held at an elevated temperature. Matrix materials are deposited on the substrate structure via a standard CVD reaction. Because of such advantages. CVI is vi/idely applied for the fabrications of fiber/whisker reinforced/ceramic matrix composites, including SiC/SiC composites.*""^ The SiC nanov^-ires in our previous SNV-TY were directly grown on a plain-woven Tyranno-SA fabric sheet""^ using a typical CVI system. Here we report, for the first time, the fabrication of a SiC nanowire/[iher reinforced/SiC matrix composite through the in silu growth of SiC nanowires in the composite preform before the CVI-matrix intiltration. A thin carbon coating was deposited on the nanowires to modify nanuwire/matrix interfacial bonding. The main interests were the flexurat strength and fracture toughness of the composite and the effects of the SiC nanowires. II. Experimental Procedures (I) Processes for Carbon-Coaled SiC Nanowires and SiC/SiC Composites The nanowires were synthesized directly in a composite preform using a CVI system, as schematically shown in Fig. 1. The composite preform was prepared by stacking 14 sheets (40 mm in diameter) of plain-woven Tyranno SA fiber cloths in a 0790° \ geometry. The stack was located in the vertical hot chamber of the CVI furnace, as shown in Fig. 1. Before the nanowire-growing process was begun, the furnace (preform) was heated to 1223 K under vacuum (<0.1 Pa) using a rotary pump. The rotary pump then was replaced by a diaphragm pump, and CH4 was introduced into the specimen chamber at a constant flow rate t)f 200 seem (standard cubic centimeters per minute) to deposit a carbon coating on the fibers in the preform as the fiber/matri\ intcrlayer. The exhausted gas was evacuated using the diaphragm pump. During the CVD carbon-coat ing process, the pressure in the chamber was maintained at 14.7 kPa (by adjusting the evacuation rate of the diaphragm pump). After the specimen was processed for 70 min. the CH4 flow was stopped, and the system was maintained at 1223 K for 10 min to make certain of complete reaction of the residual CH4 in the chamber. The specimen then was heated to 1373 K and MTS was introduced into the chamber to grow SiC nanowires in the preform. The MTS was carried by hydrogen-gas using a typical bubble system (Fig. 1). A bypass hydrogen flow also was used to dilute the MTS concentration in the chamber. During this process, the MTS bath temperature was maintained at 293 K. The flow rates of carrier and bypass hydrogen-gas were 20 and 1000 seem, respectively. Therefore, the flow rate of the MTS was 1.8 X lO""" mol/min. The nanowire process was conducted for 2 h at a decreased specimen chamber pressure of 1.47 kPa. The MTS and hydrogen-gas flows then were stopped. After the system was maintained at this level for—10 min, it was cooled to room temperature, and the weight of the preform ft-as measured to estimate the amount of SiC nanowires in the preform. One fabric sheet from the cenier of the preform was taken out for microscopic examination before the synthesized SiC nanowires in the preform were CVD carbon coated again. The carbon coating conditions were the same as described before but for a shorter processing time, i.e., 20 min. This carbon coating was supposed to decrease the bonding strength between the nanowires and the matrix. After this carbon-coat ing process, another fabric sheet was used to examine the carbon coating on the nanowires. After the completion of the second carbon-coating process, the preform was compressed with a set of graphite fixtures to 2.2 mm in thickness, which resulted in a volume fraction of fibers of -^43%. Finally, the preform was densified with SiC matrix material for 18 h using a CVI process^' at 1273 K under 14.7 kPa. The total hydrogen flow was 1000 seem, and the MTS:hydrogen volume ratio was 1:10. The composite was named NF-C (nanowire- and fiber-reinforced composite). A conventional Tyranno-SA/SiC composite, named F-C. also was fabricated for comparison. A similar CVI process as described above was used, except the nanowire portion of the process was deleted. Fig. 2. SiC nanowires on the Tyranno-SA fabric sheet
1722 Journal of the American Ceramic Sociery-Yang et al Vol, 87. No 9 (2) Flexure and Toughness Tests 800 direction, which is confirmed using the defocus SAED technique The faint streaks along the nanowire axis in the SAEd are due to the dense stacking faults along the direction. Figure 3(b) shows that the B-SiC nanowire is covered with a continuous + Cu peaks, from TEM copper mesh amorphous coating with a thickness of -5 nm. This amorphous coating did not appear at the surfaces of the Sic nanowires before the CVD carbon-coating process. As mentioned before, during the CVD carbon-coating process, only CH, was used as the source s. Therefore, this am us layer is concluded to be carbon 200 coating from the thermal decomposition of CHa. This has been … SiC nanowire confirmed using comparative EDS studies, as shown in Fig. 3(c) The relative intensity of carbon in the Sic nanowires increased after they were processed with the carbon-coating deposition process (the carbon: silicon atomic ratio increased from near toichiometric to-60/40). The copper peaks in the spectrum were from the copper mesh on which the nanowires were loaded. When Energy /kev the same process was used, the carbon coating was frequently Fig. 3. (a) and(b) HRTEM images with insert in (a) of SAED and deposited on continuous SiC fibers as the fiber/matrix interlayer in EDS spectrum of an individual carbon-coated SiC nanowire
1722 Journal of the American Ceramic Society—Yang et al. Vol. 87. No. 9 (2) Flexure and Toughness Tests The flexural strength and fracture behaviors were investigated using three-point bending tests. Rectangular specimens were cut from the composites parallel to one of the fiber bundle directions and were carefully polished to dimensions of 30 mm X 4.0 mm X 1,5 mm. The measurement of fracture toughness was conducted using single-edge notched-beam (SENB) specimens of 25 mm x 3.0 mm X 2,0 mm. The precracks were cut using a diamond blade 0.3 mm in width and 1.4 mm in depth. The bridging distance was 16 mm. and the loading velocity was 0.5 mm/min for flexure and SENB tests. The flexural strength was derived from the load/ displacement curves of three bending specimens, according to ASTM C 1341-97.^^ The fracture toughness (KiJ was calculated from load/displacement curves of four SENB tests according to ASTMC 1421-99." (3) Microstructure Characterization The nanowires were examined using scanning electron microscopy (SEM; Model JSM-6700F. JEOL. Tokyo, Japan), highresolution transmission electron microscopy (HRTEM; Model JEM-3000F. JEOL), and selected-area electron diffractometry (SAED; Model JEM-3000F, JEOL). The chemistry of the nanowires was examined using energy-dispersive spectroscopy (EDS) attached to the HRTEM. The interlayers, microstructures of the composite, and fracture surfaces after tests were examined using SEM. i n. Results and Discussion (1) SiC Nanowires on the Fabric Sheet Figure 2(a) shows the SEM image of a Tyranno-SA fabric sheet after the nanowire growth process. The sheet is covered with a spongelike layer, which is an aggregation of nanowires, as confinned by the higher-magnification SEM image in Fig. 2(b), The nanowires are generally several tens to more then one hundred micrometers in length and randomly oriented with straight or curved morphologies. The diameter of most of the nanowires is — 100 nm. However, nanowires with diameters direction, which is confirmed using the defocus SAED technique. The faint streaks along the nanowire axis in the SAED are due to the dense stacking faults along the direction. Figure 3(b) shows that the P-SiC nanowire is covered with a continuous amorphous coating with a thickness of ~5 nm. This amorphous coating did not appear at the surfaces of the SiC nanowires before the CVD carbon-coating process. As mentioned before, during the CVD carbon-coating process, only CH4 was used as the source gas. Therefore, this amorphous layer is concluded to be carbon coating from the thermal decomposition of CH4. This has been confirmed using comparative EDS studies, as shown in Fig, 3(c). The relative intensity of carbon in the SiC nanowires increased after they were processed with the carbon-coating deposition process (the carbon:silicon atomic ratio increased from near stoichiometric to —60/40). The copper peaks in the spectrum were from the copper mesh on which the nanowires were loaded. When the same process was used, the carbon coating was frequently deposited on continuous SiC fibers as the fiber/matrix interlayer in (c) -coate(J SiC nanowire XJ:^ Cu peaks, from JWl copper mesh • —-SiC nanowire J 'L. 2 4 6 Energy / keV 10 Fig. 3. (a) ana (b) HRTEM images with insert In (a) (if SAED mid (c) EDS spectrum 0!' an individual carbon-coated SiC nanowire
September 2004 Process and Mechanical Properties of in Situ SiC Nanowire-Reinfarced CVI SiC/SiC Composite 1723 Table L. Density, Interlayer, and Mechanical Properties of the Composites Carbon laver ssm) U (MPa '(MPw/(Mg/m'n K,(MPam) Kl (MPam A(Mg/n') NFC262±0.03 660±77 250±29 16.3±3.3 6.2±1.2 F-C 2.74±0.04 590±78 210±28 15.2±3.5 5.5±1.3 and Ai, are normalized ultimate flexural strength and fracture toughness, respectively, using composite densities. SiC/SiC composites. Many other SEM and TEM examinations features at their fracture surfaces, as shown typically in Figs. 5(a) confirmed successful growth of SiC nanowires directly in the and(b) for composite NF-C and Figs. 5(c)and (d) for composite composite preform with in situ deposition of thin and uniform carbon coating on all the nanowires, although the amount of the esults in sound fiber pullout fracture behaviors. Composite F-C nanowires was relatively small. exhibits slightly longer fiber pullouts than composite NF-C. likely because of its thicker carbon interlayer (Table I), which produces (2) Fabricated Composites Table I shows the density and the thickness of carbon fiber/ No clear debonding and pullout of the SiC nanowires from the hatrix interlayers of the two composites. The density is the matrix are observed at the fracture surfaces of composite NF-C. average value of the seven rectangular specimens of each com except occasional observations of fragments of nanorods posite from their masses and volumes. Composite NF-C has an average density of 2.62=0.03 Mg/m, which corresponds to a (4) Effect of SiC Nanowires porosity of -13%. A 60 nm carbon coating was deposited on the fibers as the fiber/matrix interlayer, as shown in Fig. 4(a) and the As shown in Table I. the two composites possess different nserted high-magnification SEM image. in which the fiber carbon interlayer and average density, which are important factors matrix. and the carbon coating around the fibers are evident. that affect mechanical performance. These differences make it However, no SiC nanowires are noted in the image. Similarly, at difficult to clearly understand the effects of the nanowires on he cross section of the composite where the matrix is densel mechanical properties by simply comparing the strength data of deposited. it is difticult to observe SiC nanowires in the matrix or the two composites. In previous studies . tensile and flexural around the fibers using other SEM examinations. Microscale trengths of a series of plain-woven Tyranno SA fiber-reinforced CVI-SiC/SiC composites showed close carbon layer thickness rods/structures are readily observed in areas with an insufficiently deposited matrix as shown typically in Fig. 4(b). Such microscale dependence up to-100 nm. When the carbon layer thickness ncreased from 50 to 100 nm. the ultimate tensile and flexural structures are not found in the companion composite, F-C. These strengths increased from 410+ 92 and 189+27 MPa to 606+ microscale structures are believed to have developed from the original nanowires during the CVI matrix densification process 28 and 285+ 20 MPa, respectively. The thickness of the carbon that followed. The absence of SiC nanowires in the dense matrix interlayer in composite NF-C is 60 nm, and it is 120 nm in in NF-C is likely because of the very thin carbon coating (--5 nm). composite F-C. Therefore, the real contributions of the SiC which is not sufficient to form a clearly visible interface between nanowires on the strength and toughness of composite NF-C might the nanowires and the matrix under current SEM examinations be partially covered by the effects of the thickness of the carbon Composite F-C has a thicker carbon interlayer, 120 nm, and high interlayer. In addition, statistical studies" have shown that the average density, 2.74 0.04 Mg/m flexural strength of SiC/SiC composites also increase with higher composite densities. However. composite NF-C shows higher (3) Flexural Strength and Fracture Toughness flexural strength and fracture toughness compared with composite F-C, despite its thinner carbon fiber/matrix interlayer and lower The obtained ultimate flexural strength(o )and fracture tough- density. The increased strength and toughness are believed to be ness(Ki)are summarized in Table L. The composite NF-C shows due to the contribution of the incorporated SiC nanowires. Here, average flexural strength and fracture toughness. 660+77 MP when the average strength and fracture toughness of composit and 16.3+3.3 MPam". respectively, which are slightly higher densities are simply normalized, composite NF-C shows% than those of composite F-C. The two composites have similar and -13% increases in flexural strength and fracture toughness Matrix Fiber Fiber Bilayer 5 m k Matrix Micro-rods 100m 5 u m Fig. 4. SEM images of the (a) cross section and (b) microrods/structures in composite NF-C
September 2004 Process iiinl Mechanical Properties of in Situ SiC Nanmvirc-Reinforced CVI SiC/SiC Composite Table I. Density. Interlayer. and Mechanical Propt'rties of the Composites Composite Dcnsii Carbon layer thickness Iniii) • r,, lMP;i) (MP;i-m' NF-C F-C 2.62 ± 0.03 2.74 ± 0.04 60 660 ± 77 120 590 ±7 8 250 ± 29 210 ± 28 16.3 ±3.3 15.2 ±3.5 iMPii-m"- 6.2 + 1.2 5.5 ± 1.3 ir;, and A*;, arc normalised ultimate flexural stroiiyih ami Iniilurc touj^hness, respociively. u.'.ing composite dcnsitit 1723 SiC/SiC composites.*''' Many other SEM and TEM exaniin;iti(ins confirmed successful growth of SiC nanuwires directly in the composite preform with in situ dcpositit)n of ihin and Linilbrm carbon coating on all the nanowires. although the amount of the nanowires w;is relatively small. i2) Fabricated Composites Table I shows the density and the thickness of carbon tlber/ matrix interlayers of the iwo ctimposites. The density is the average value of the seven rectangular specimens of each composite from their masses and volumes. Composite NF-C has an average density of 2.62 ± 0.03 Mg/tn"*. which corresponds lo a porosity of— 13%. A 60 nm carbon coating was deposited on the fibers as the tlber/inairix interlayer. as shown in Fig. 4(a) and the inserted high-magnillcation SEM image, in which the fiber, matrix, and ihe carbon coating around the fibers are evidcni. However, no SiC nanowires ari? noted in the image. Similarly, al ihe cross section of the composite where the matrix is densely deposited, it is difficult to observe SiC nanowires in ihe matrix or around the fibers using other SEM examinations. Microscale rods/struclures are readily observed in areas with an insufficiently deposited matrix, as shown typically in Fig. 4(b}. Such microscale structures are not found in the companion composite. F-C. These microscale structures are believed to have developed from ihe original nanowires during the CVI matrix densitlcatioii process that followed. The absence of SiC nanowires in the dense matrix in NF-C is likely because of the very thin carbon coating (—5 nm). which is not sufficient to form a clearly visible interface between the nanowires and the matrix under current SEM examinations. Composite F-C has a thicker carbon interlayer. 120nm, and higher average density, 2.74 ± 0.04 ' (3) Flexural Strength and Fracture Toughness The obtained uhimate tlexural strength (uj and fracture toughness (A",,,) are summari/cd in Table I. The composite NF-C shows average tlexural strength and fracture toughness. 660 ± 77 MPa and 16.3 ± 3.3 MPa-m"-. rcspeciiveiy. which are slightly higher than those of composite F-C, The two composites have similar features al their fracture surfaces, as shown typically in Figs. 5(a) and |b) for composite NF-C and Figs. 5(c) and (d) for composite F-C. Inierfacial debonding occurs al the fiber surfaces, which results in sound fiber pulknit fracture behaviors. Composite F-C exhibits slightly longer fiber pullouts than composite NF-C. likely because of its thicker carbon interlayer (Table I), which produces lower interfacial shear strength/friclional stress in the materials.'' No clear debonding and pulloui oi' the SiC nanowires from the matrix are observed at the fracture surfaces of composite NF-C. excepi occasional observations of fragments of nant>rods. (4) Effect of SiC Nanowires As shown in Table 1. the iwo composites possess different carbon interlayer and average density, which are important factors that affect mechanical performance. These differences make it difficult to clearly understand the effects of the nanowires on mechanical properties by simply comparing the strength data of the two composites, in previous studies.'''' tensile and flexural sirengths of a series of plain-woven Tyranno-SA llber-reinforced CVI-SiC/SiC composites showed close carbon layer thickness dependence up to ~IOO nm. When the carbon layer thickness increased from 50 lo 100 nm. the ultimate tensile and tlexural strengths increased from 410 ± 92 and 189 ± 27 MPa to 606 ± 28 and 285 ± 20 MPa. respectively. The thickness of the carbon interlayer in composite NF-C is 60 nm. and it is 120 nm in composite F-C. Therefore, the real contributions of the SiC nanowires on the strength and toughness of composite NF-C might be partially covered by the etfects of the thickness of the carbon interlayer. In addition, statistical studies"'* have shown that Ihe Hexural strength of SiC/SiC composiics also increase with higher composite densities. However, composite NF-C shows higher tlexural strength and fracture toughness compared with composite F-C, despite its thinner carbon tlber/matrix interlayer and lower density. The increased strength and toughness arc believed to be due to the contribution of the incorporated SiC nanowires. Here. when the average strength and fracture toughness of composile densities are simply normali/.ed. composite NF-C shows —20^*^ and —13'^ increases in tlexural strength and fracture toughness Fig, 4. SHM image s of ihc fa) cross ii and ib) microrodsAirtictures in conipusiie NF-C
Journal of the American Ceramic Sociery-Yang et al. Vol 87. No 9 100am 5 u m 100m 10m Fig. 5. SEM images of fracture surfaces and fiber pullouts of the composites after SENB tests(a) and (b) of composite NF-C and (c)and (d) of composite Table I) respectively. over composite F-C. After the mechanical direction in composite NF-C. These bundles were not fully tests were completed. some of the specimens were carefully densified with matrix: therefore, the microrods/structures within side-surface polished to examine the transverse matrix cracking them(as mentioned before) are clearly seen. These microscale c haviors. Figure 6 shows typical fractured fiber bundles perpen structures provide connections between the individual fibers in dicular (Fig. 6(a)and parallel (Fig. 6(b) to the main stress these poorly densified fiber bundles, which results in some (a (b) 20 u m 10m Fig. 6, Microscale structures in porous fiber bundles in composite NF-C after failure
1724 Jdiinuil of the American Ceramic Socteiy—Yanf- el al. Vol. 87. No. 9 Fig. 5, F-C). S EM inuge s of fracture sutfaces and fiber pullouUs of the composifes after SENB tests ((a) and (b) ol" composite NF-C and (cl and (d) of composile (Table I), respectively, over composite F-C. After the mechanical lests were completed, some of the specimens were carefully side-surface polished to examine the (ransverse matrix cracking hehaviors. Figure 6 shows typical fractured fiher bundles perpendicular (Fig. 6(a)) and parallel (Fig. 6(b)) to the main stress direction in composite NF-C. These bundles were not fully densified with matrix; therefore, the microrods/struclures within them (as mentioned before) are clearly seen. These microscale structures provide connections between the individual fibers in these poorly densified fiber bundles, which results in some Fig. 6. Microscale structures iti porotis tlber bundles in composite NF-C after failure
September 2004 Process and Mechanical Properties of in Situ SiC Nanowire Reinforced CVI SiC/SiC Composite load-transferring bridges between them. The SEM images in Fig. provides more reasonable nanowire/matrix interfacial bonding 6 show that some of the microrods are broken while others remain and/or increased load of the nanowires unbroken. More fracture energy is dissipated when the matrix cracks propagate through these bundles to break the microrods. On References the other hand. the unbroken microstructures enable the cracked matrix to suspend a certain level of strength. It is believed that D Brewer, "HSR/EPM Combustor Materials Development Program. "Mater. Sci. these SiC-nanowire- based microscale structures, as well as these Eng,A261.28491(19 and O.J. Schwarz,"Advanced SiC Composites for Fusion Applica SiC nanowires buried in densely deposited matrix, serve as tion. " J. Nucl. Mater. 219 3-14(1995) H. Matsui, S. Jitsuknwn, and S Matsuda, "Interactions Between Fusion Materials R&D and Other Technologies. However, when the very high strength of the single-crystal SiC J Nucl. Mater. 283-287, 20-27(2000) nanowires is considered, o the reinforcement effects of the sic C: Droillard and J, Lamon, "Fracture Toughness of 2-D Woven SiC/SiC nanowires in the current composite are not significant. One of the Cv- Composites with Multi-Interlayered Interlayers"/.Am. Ceram. Soc.79 141 important issues for fiber or nanowire reinforcement of materials is tw. Yang. A Kohyama, Y, Katoh, H. Araki. J. Yu, and T Noda, "Effect of C and the control of the interfacial bonding between the reinforcement SiC/C Interlayers on Mechanical Behavior of Tyranno-SA Fiber-Reinforced Sic and the matrix, which must be neither too strong nor too weak ome degree of fiber/nanowire pullout allows energy to be ab- R. Naslain, "Fiber-Matrix Interphases and Interfaces in Ceramic-Matrix Compos- ites Processed by CVI, "Compos. Interfaces, 1 341 253-86(1993) orbed in breaking reinforcement/matrix bonding. The fracture I M. Besmann, D. P, Stinton, E R. Kupp S Shanmugham, and P. KLiaw surface examinations reveal that the bonding between the Tyranno-SA fibers and the matrix in composite NF-C is reason- 45%.- 9 g/rfaces n Ceramic Compesite Mater. Res Se, swnp. Pre. omposite, No debonding and pullout of SiC nanowires from(s able, because sound fiber pullouts occur during the failure of the w. Yang, T Noda, H. Araki. J. Yu, and A Kohyama."Mechanical Properties of Several Advanced Tyranno- SA Fiber-Reinforced CvI-SiC- Matrix Composites, Ala- ter.Sdi.Eng,A345.28-35(2003 matrix is observed at the fracture surfaces, which indicates a too IE W. Wong, P. E Sheehan, and C M. Lieber. "Nanobeam Mechanics: Elastici strong bonding between the nanowires and matrix. This is likely 277 126)1971-75(1997) due to the very thin(-5 nm) compliant carbon coating on the W. R. L Lambrecht, B. Segall, M Methfessel, and M. van Schilfgaard nanowires, Such a thin carbon coating is insufficient to produce "Calculated Elastic Constants and Defonnation Potential of Cuhic SiC, Phvs. Rev ound bonding strength between the nanowires and the matrix to Condens Matter, 44, 3685-91 11991) allow interfacial debonding during the failure of the composite IP. Calvert, "A Receipe for Strength. "Nature (London), 399. 210-11(1999) IR. W. Sirgel, S K Chang,, B J. Ash. J. Stone. P. M. Ajayan R W.Doremus, and Bridging of matrix cracks by the Sic nanowires rarely takes place L S Schadler, "Mechanical Behavior of Polymer and Ceramic-Matrix Nanocompos- in this case. Therefore, the efficiency of the reinforcement of SiC tes,"Scr.Maer,44.2061-64(2001 E. Flahaut. A. Peigney, Ch. Laurent, Ch. Marliere, F, Chastel, and A. Rousset. nanowires might be largely decreased. Nevertheless, this smdy"Carbon Nanotuhe-Metal Oxide Nanocomposites: Microstructure,Electrical Conduc- demonstrates the possibility to fabricate in situ SiC-nanowire- einforced ceramic-matrix composites with hopefully markedI D. Zhan, J Huntz, J. Wan, and A. Mukherjee,"Single. Wall Carbon Nano- improved strength and toughness, which provides thicker carbon ubes as Attractive Toughening Agents in Alumina- Based Nanocomposites, "Nar or alternative compliant coatings on the surfaces of the nanowires nd/or increases the load of the nanowires "T, Ishikawa, Y, Kohtoku, K Kumagawa, T. Yamamura, and T, Nagasawa gh-Strength Alkali-Resistance Sintered SiC Fiber Stable to 2200C,Nature london)39166691773-751998 H. Dai. E. W. Wong. Y, Z Lu S. Fan, and C. M. Lieber. ""Synthesis and Clrae Nature .3751291769-72(1995 IV. Conclusion I. T. Zhou N, Wang. F. C. K. Au, H L Lai, H. Y, Peng, L Bello, C.S. Lee, and S. T. Lee, "Growth and Emussion Properties of B-SiC Nanorods. " Mater. Sci. Eng A SiC-nanowire/Tyranno-SA-fiber-reinforced SiC/SiC com- A286.119-24(2000 posite was fabricated via the simple in situ growth of SiC C. C. Chiu. S. B. Desu, and C. Y Tsai."Low-Pressure Chemical Vapor nanowires directly in the fibrous preform before CVI matrix Deposition (LPCVD) of B-SiC on Si( 100) Using MTS in a Hot Wall Reactor. J. Mater.Res.,8o2617-25(1993 densification. The volume fraction of the SiC nanowires in the composite was estimated to be -5% The nanowires were characterized as pure and single-crystal H. P, Kirchner and P. Knoll, "Silicon Carbide Whiskers, "J. Am. Ceran Soc., 46 299-300(1963 B-phase SiC with diameters of several tens to one hundred nanometers and lengths of several tens of micrometers, A uniform e2D. Lespiaux. F. Langlais,R. Naslain. A. Schamm, and J.Sevely "Correlation carbon coating of-5 nm was successfully deposited on the Characteristics of Silicon Carbide Deposited from Si-C-H-CI System on Silica nanowires in situ as the nanowire/matrix interlayer Substrate. "J. Mater. Sci, 301. 1500-10(199 ow. Yang. H. Araki, A. Kohyama, Q-LHu, H. Suzuki, andT Noda, "Grow The composites with Sit nanowire showed higher flexural SiC Nanowires SiC Nanowires on Tyranno-SA SiC Fibers, J. An. Cera. Soc. $7141 733- strength and fracture toughness compared with those of the companion conventional composite, despite its thinner carbon T. M. Besmann, B. w. Sheldon, R. A Lowden, and D P. Stinton. "Vapor-Phase fiber/matrix interlayer and lower density. The increased strengt Fabrication and Properties of Continuous-Filament Ceramic Composites, Science Washington, DC.253.104-10901991) and toughness were attributed to the SiC nanowires in the R. Naslain and F. Langlais, "CVD- Processing of Ceramic-Ceramic Composite mposite. However, the increased strength and toughness were As in Tailoring Multiphase and Composite Ceramics. Edited by R.E. not remarkable, and little debonding and pullout of SiC nanowires Tressler, G. L Messing. G. G. Pantano, and R. E Newnham. Plenum, New York, observed. It seemed that the carbon coating on the current sic Advanced Ceramic Composites. ASTM Designation C 1341-97 2000 ASTM Annual nanowires was not thick enough to produce a sound interfacial Book of Standards, Part 15. ASTM International, West Conshohocken. PA bounding to allow nanowire/matrix debonding during the failure of H Standard Test Method for Determination of Fracture Toughness of Advanced the composite. The efficiency of the reinforcement by the SiC Annual Book of Standands Part 15. ASTM International, West Conshohocken.PA nanowires was, therefore, believed to be decreased The present process demonstrated the possibility to fabricate in situ SiC-nanowire-reinforced ceramic-matrix composite with Transactions, Vol. 144. Advanced SiC/SiC Ceramie Composites-Developmernts and Applications to Energy Systems. Edited by A Koby ama, M. Singh, H. T. Lin, and Y hopefully markedly improved strength and toughness, which Katob. American Ceramie Society. Westerville. OH, 2002
September 2004 Process and Mechanical Properties of in Situ SiC Nanowire-Reinforced CVl SiC/SiC Composite 1725 load-transferring bridges between them. The SEM images in Fig. 6 show thai some of the microrods are broken while others remain unbroken. More fracture energy is dissipated when the matrix eracks propagate through these bundles to break the microrods. On the other hand, the unbroken microstructures enable the cracked matrix to suspend a certain level of strength. It is believed thai these SiC-nanowire-based microscale structures, as well as these SiC nanowires buried in densely deposited matrix, serve as additional reinforcement elements besides the Tyranno-SA fibers, which results in a contribution to the increase of flexural strength and fracture toughness. However, when the very high strength of the single-crystal SiC nanowires is considered,"^ the reinforcement effects of the SiC nanowires in the current composite are not significant. One of the important issues for fiber or nanowire reinforcement of materials is the control of the interfacial bonding between the reinforcement and the matrix, which must be neither too strong nor too weak; some degree of fiber/nanowire pullout allows energy to be absorbed in breaking reinforcement/matrix bonding. The fracture surface examinations reveal that the bonding between the Tyranno-SA fibers and the matrix in composite NF-C is reasonable, because sound fiber pullouts t>ccur during the failure of the composite. No debonding and pullout of SiC nanowires from the matrix is observed at the fracture surfaces, which indicates a too strong bonding between the nanowires and matrix. This is likely due to the very thin (~5 nm) compliant carbon coating on the nanowires. Such a thin carbon coating is insufficient to produce sound bonding strength between the nanowires and the matrix to allow interfacial debonding during the failure of the composite. Bridging of matrix cracks by the SiC nanowires rarely takes place in this case. Therefore, the efficiency of the reinforcement of SiC nanowires might be largely decreased. Nevertheless, this study demonstrates the possibility to fabricate in situ SiC-nanowirereinforced ceramic-matrix composites with hopefully markedly improved strength and toughness, which provides thicker carbon or alternative compliant coatings on the surfaces of ihe nanowires and/or increases the load of the nanowires. IV. Conclusion A SiC-nanowire/Tyranno-SA-fiber-reinforced SiC/SiC composite was fabricated via the simple in .situ growth of SiC nanowires directly in the fibrous preform before CVI matrix densification. The volume fraction of the SiC nanowires in the composite was estimated to be —5%. The nanowires were characterized as pure and single-crystal P-phase SiC with diameters of several tens to one hundred nanometers and lengths of several tens of micrometers. A uniform carbon coating of ~5 nm was successfully deposited on the nanowires in situ as the nanowire/matrix interlayer. The composites with SiC nanowires showed higher flexural strength and fracture toughness compared with those of the companion conventional composite, despite its thinner carbon llber/matrix interlayer and lower density. The increased strength and toughness were attributed to the SiC nanowires in the composite. However, the increased strength and toughness were not remarkable, and little debonding and pullout of SiC nanowires from the matrix at the fracture surface of the composite were observed. It seemed that the carbon coating on the current SiC nanowires was not thick enough to produce a sound interfacial bounding to allow nanowire/matrix debonding during the failure of the compt>site. The efficiency of the reinforcement by the SiC nanowires was, therefore, believed to be decreased. 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