- Email: [email protected]
Carbon fiber reinforced silicon carbide mini-composites-solution approach
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
journal homepage: www.elsevier.com/locate/jmatprotec
Carbon fiber reinforced silicon carbide mini-composites-solution approach N. Padmavathi a,b,∗ , J. Subrahmanyam a , K.K. Ray b , R. Mohanrao a , P. Ghosal a , Sweety Kumari a a b
Defence Metallurgical Research Laboratory, Kanchanbagh and Hyderabad 58, India Metallurgical and Materials Science Engineering, Indian Institution of Technology, Kharagpur 721302, India
a r t i c l e
i n f o
a b s t r a c t
Article history:
Carbon fiber reinforced composites are widely used as aerospace materials owing to their
Received 4 July 2007
superior properties like high specific strength and high fracture toughness. Oxidation above
Received in revised form
600 ◦ C is one of the major problems encountered with carbon–carbon composites. Replace-
1 November 2007
ment of carbon matrix with SiC is one of the methods to improve the oxidation resistance.
Accepted 12 November 2007
A relatively inexpensive process is developed for the fabrication of carbon fiber reinforced SiC composites through solution approach. This process allowed forming high purity and fine-grained matrix in the fiber reinforced composites. Carbon fiber tow mini-composites
Keywords:
were fabricated using this process. Typical tensile strength and fracture energy of these com-
Carbon fiber
posites were 316 MPa and 0.74 MJ/m3 at room temperature (RT). The composites exhibited
SiC
load-carrying capability after crack initiation.
Solution approach
© 2007 Elsevier B.V. All rights reserved.
Composite Tensile strength
1.
Introduction
Fiber-reinforced-ceramic-matrix composites have recently attracted attention for use in high-temperature structural applications (Stinton et al., 1986; Prewo, 1989). The primary reason for this interest lies in the assumption that strong ceramic fibers can prevent catastrophic brittle failure in ceramics by providing various energy-dissipation processes during crack advance (Evans, 1989). Early works utilizing ceramic fibers in ceramic matrices demonstrated the potential of this approach (Prewo and Brennan, 1980; Stinton et al., 1985, 1986; Shetty et al., 1985; Naslain et al., 2001; Tressler, 1986; Fitzer and Gadow, 1986). Carbon fiber reinforced silicon carbide (C/SiC) composites provide excellent thermo-mechanical properties at tempera-
tures up to 2000 ◦ C (Zhong et al., 1998; He et al., 2001; Mentz and Muller, 2006). Their high strength-to-weight ratio, which sustains at the high temperatures, makes them potential candidates for highly demanding engineering applications such as heat shields and structural components for re-entry space vehicles, high performance brake discs, and ultra-high temperature heat exchanger tubes. For these elements, C/SiC composites have the most favourable properties and are, therefore, at the challenging front of the contemporary materials research. However, a generic problem that must be overcome is that ceramic fabrication processes tend to mechanically and chemically damage the fibers when they are consolidated within a ceramic matrix. For example, the fibers may be broken by a pressing operation, or the high sintering temperature required
∗ Corresponding author at: Defence Metallurgical and Research Laboratory, Kanchanbagh and Hyderabad 58, India. Tel.: +91 9440413314; fax: +91 40 24340683. E-mail address: nandigam [email protected] (N. Padmavathi). 0924-0136/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.11.124
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
to densify the ceramic matrix may damage the fibers or cause them to react chemically with the matrix. C/SiC composites can be fabricated by a CVI infiltration process, yet their high performance is attained at prohibiting costs for many applications (Kim et al., 1993; Kenji et al., 2001; Cheng et al., 2001). Polymer impregnation and pyrolysis (PIP) derived SiC matrix contains free carbon and oxygen (Sato et al., 1999; Boisvert and Diefendorf, 1988). Liquid silicon infiltration processes also a common fabrication route, offer some cost advantage, but produces SiC matrix along with free silicon (Krenkel, 2001). New processing techniques are to be developed to get high quality SiC matrix at low cost. Soft solution processing described in this paper allows fabrication of advanced fiber reinforced composites using aqueous solutions, without expensive equipment, providing an environmental friendly route. In this study, carbon fiber reinforced SiC composites have been prepared by a solution approach (Weimer et al., 1993). Solution approach or soft chemistry route involves the low temperature modification of existing solid structures to form new solids that retains many of the structural features of the precursor phase. A mixture of colloidal silica and sucrose in solution form were used as a precursor for formation of SiC matrix. Drying and heating to 500 ◦ C converts the sucrose into carbon and SiC forms by the carbothermic reduction of silica at high temperatures by the following reaction, SiO2 + 3C → SiC + 2CO
(1)
Mini-composites were fabricated by impregnating the solution phase into a carbon fiber tow, drying, converting sucrose into carbon phase and finally reducing silica with carbon to form SiC. The principal advantage of this approach is the relatively inexpensive ingredients that constitute the solution
435
phase, to get fine-grained SiC matrix in a fiber reinforced composite (Subrahmanyam et al., 2004).
2.
Experimental procedure
2.1.
Raw materials
The carbon fiber used in the study was supplied by Tae Kwang Ind. Co. Ltd and Japan, Grade-TZ-300 with 12,000 filaments and filament diameter of 6 m (12 K). Colloidal silica was obtained from Bee Chem. Chemicals Company, Kanpur and India, and Sucrose from Qualigens Fine chemicals, India.
2.2.
Composite-synthesis
Grade-TZ-300 carbon fiber tow was initially washed with acetone to remove impurities and sizing. The composition of the precursor solution is 0.25 M of colloidal silica and 0.876 M of sucrose in distilled water. The fiber tows were vacuum impregnated with precursor solutions of SiC (20–25 ml per 80 mg of fiber tow) to obtain mini-composites with SiC matrix. Impregnated tows were dried carefully at room temperature and at 60–70 ◦ C for 12 h. Dried tows were carbonized at 500 ◦ C for 1 h for converting sucrose into carbon. These steps were repeated for five to six cycles. After the desired impregnation/heat treatment cycles, carbonized tows were pyrolysed at 1600 and 1700 ◦ C for 3 h under flowing argon to obtain the SiC matrix. Flow sheet for the process is shown in Fig. 1. Bulk densities of the samples were measured with deionised water as immersion medium according to the Archimedes principle. The densities of the composites increased with the number of infiltration cycles. Further cycles were not recommended after completing of six cycles due to
Fig. 1 – Flow sheet for the process.
436
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
the rapid decreasing effectiveness of further infiltration. Density of the composite was 1940 kg cm−3 and porosity is 30%. The volume fraction of fiber and matrix are 88% and 12%, respectively.
2.3.
Microstructural characterization
Philips Model No. PW 320 is used to carry out X-ray diffraction (XRD) of samples to identify the crystalline phases of the powder samples produced from the precursor solution as well as the phases from mini-composites. The microstructures of the mini-composites were examined using an optical microscope as well as a scanning electron microscopy (Model No. 7060, SEM- Oxford Leo 440i) equipped with an energy dispersive X-ray spectrometer (EDS).
2.4.
Evaluation of mechanical properties
Relatively simple bend testing (Flexure testing) is generally used for ceramic samples. However it is not suitable for ceramic matrix composites due to non-uniform stress state and brittle nature of composites matrix. In the present work mini-composite approach, single or a few tows are combined together and impregnated to get the composite (Naslain et al., 1999; Lamon et al., 1995). The benefits with the minicomposite approach are better utilization of the expensive fiber material and the mini-tow composite tensile sample has a uni-axial tensile stress condition, which gives the more appropriate data, for process development compared to flexure testing and it allows study of various process variables in a shorter time (Gonczy et al., 1997; Larsen and Stuchly, 1990). The mechanical tests were performed using a servohydraulic testing machine (Model 8801; Instron Corp., UK) at room temperature under a displacement rate of 0.5 mm/min. The load was measured using a 1000 N load cell. Specimens are prepared by following the procedure outlined for minicomposites at room temperature with a gauge length of 50 mm and using paper tabs for gripping with a suitable adhesive (Naslain et al., 1999). Once the mini-composite sample was mounted on the machine, the cardboard tab supporting the fiber was cut and testing carried out. The properties of mini-composites are characterized by two separate mechanical parameters in this study: (1) tensile strength was calculated from the stress–strain curve of each composite and (2) fracture energy was measured by area under the curve for each composite. Fracture energy gives the damage tolerance of the composite. Weibull modulus (m) reflects the degree of variability in the strength values; where a higher number indicates less scatter in data. For ceramics, the values are in the range of 5–20 (Davidge, 1979; Lamon et al., 1993).
3.
Results and discussion
3.1.
Evaluation of the microstructures
XRD pattern of carbon fiber tow for comparison with tow composite is presented in Figs. 2 and 3 shows that the XRD pattern of the powder samples produced from the precursor
Fig. 2 – XRD pattern for carbon fiber tow.
solution, which confirms the formation of -SiC. However, the powders prepared at 1600 and 1700 ◦ C not showing any difference in terms of peak intensity. Fig. 4 shows the XRD patterns of the mini-composites, synthesized at two different pyrolysis temperatures. The C-SiC composite consists mainly -SiC and minor graphite peaks originating from the fiber. There is no trace of unreacted SiO2 in diffraction patterns. From the diffractogram (Fig. 4), it is clear that at a pyrolysis temperature of 1600 ◦ C, -SiC peaks are just appearing and at 1700 ◦ C, the peak intensity increased considerably. A graphite peak is also present in these diffractograms indicating the presence of carbon fibers. This may be due to the fact that at 1600 ◦ C, micro-crystalline -SiC forms from an amorphous solution phase and at 1700 ◦ C the size of the crystallites increases (Naslain, 2004).
Fig. 3 – XRD pattern for the powders prepared through soft solution route (a) at 1600 ◦ C and (b) at 1700 ◦ C/3 h.
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
437
Fig. 4 – XRD pattern for the C/SiC composites prepared through soft solution route pyrolysed at (a) 1600 ◦ C and (b). 1700 ◦ C for duration of 3 h.
3.2.
Microstructures of composites
Fig. 5 shows the SEM pictures of carbon fiber tow and mini-composites which is opened vertically to view the microstructure of the interior surface of the composite. It shows the distribution of matrix between the fibers through the composite. It is difficult to prepare cross-sectional metallographic samples from the mini-composites. During the preparation of the porous and brittle matrix of mini-composite
Fig. 6 – Cross-sectional electron micrographs of composite: (a) SEM micrograph and (b) optical micrograph with polarized light.
was severely damaged. During dry cutting using a diamond saw without coolant, the brittle fibers were torn and pulled out. When cut in wet condition, the removed carbon residue and fiber remnants were smeared on the surface, making further observations difficult. Finally it was possible to view the microstructure of the composites by filling the porosity with a low-viscosity resin. This fixed the individual fibers and facilitated polishing of the sample. And the charged surfaces of resin portions were carefully avoided while observing the polished section of the mini-composites. The cross section of the polished composite was observed under SEM. Fig. 6(a) shows a typical micrograph of pyrolysed CMC showing distribution of fibers and matrix. Fig. 6(b) shows the optical micrograph using polarized light of minicomposite showing individual fibers clearly separated by the SiC matrix.
3.3.
Fig. 5 – SEM micrograph of (a) carbon fiber and (b) C/SiC composites fabricated by solution approach.
Mechanical properties
Tensile tests were performed on carbon fiber tows and the mini-composites. Fig. 7(a) shows the stress–strain curve for carbon fiber tow at room temperature. The fiber tow showed gradual decrease in load after the point of maximum load, due to individual fiber breakage. Fig. 7(b) shows the tensile behavior of two composites, one pyrolysed at 1600 ◦ C and the other processed at 1700 ◦ C. The composites pyrolysed at 1600 ◦ C showed a gradual decrease in stress after the point
438
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
Fig. 7 – Stress–strain curves at room temperature (a) carbon fiber tow and in (b). (1) Pyrolysed at 1600 ◦ C and (2) At 1700 ◦ C.
Fig. 8 – Plot to obtain the Weibull modulus of 6.6 for 14 samples pyrolysed at 1600 ◦ C tested at room temperature. Fig. 9 – Plot to obtain the Weibull modulus of 3.7 for 15 samples pyrolysed at 1700 ◦ C tested at room temperature. of maximum load, similar to the fiber tow, typical of a composite failure. The mini-composites pyrolysed at 1700 ◦ C, not showing this type of behavior. Stress decreased rapidly after the point of maximum load indicative of catastrophic monolithic ceramic type failure behavior. A possible reason for this difference could be due to the difference between the matrices obtained at these two different temperatures. As shown in the XRD patterns for these composites, the SiC matrix obtained at 1600 ◦ C is in a micro-crystalline form with low Xray peak intensities. On the other hand the peak intensities have increased considerably for the SiC matrix processed at 1700 ◦ C showing the evidence for larger crystallite size. In order to quantify these results, several mini-composite samples were prepared at both temperatures and tensile tests were carried out. Weibull statistical analysis was carried out on this data to obtain the Weibull modulus. Weibull modulus gives an idea about the scatter in the data, higher the value lowers the scatter. Figs. 8 and 9 present the Weibull
plots for the composites processed at 1600 and 1700 ◦ C, respectively. Fracture energy for each of these composites is obtained from the area under the tensile curves. The average values of the fracture energy, along with the tensile strength and the Weibull modulus are presented in Table 1. As can be seen from Table 1, carbon fiber tows without any matrix shows highest strength, fracture energy and Weibull modulus. It is well known that the fiber strength cannot be realized in the composites due to (i) fiber damage during composite fabrication process and (ii) constraint imposed by the matrix on the fiber. Among the composites, the one processed at 1600 ◦ C, gave not only substantially higher tensile strength and fracture energy, but also higher Weibull modulus indicative of less scatter in the data. As already mentioned, this may be due to the micro-crystalline nature of the matrix formed at 1600 ◦ C. In addition silicon carbide can react with the residual
Table 1 – Mechanical properties of C/SiC composites Material Carbon C-SiC at 1600 ◦ C for 3 h C-SiC at 1700 ◦ C for 3 h
Tensile strength (MPa) 1146 ± 104 316 ± 48 149 ± 38
Fracture energy (MJ/m3 ) 10.4 0.74 0.07
Weibull modulus (m) 10.5 6.6 3.7
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
silica in the matrix at high temperature resulting in weaker matrix (Clegg, 2000).
4.
Conclusions
This investigation has clearly demonstrated the possibility of fabrication of the carbon fiber reinforced SiC mini-composites by soft solution approach. X ray diffraction of the powders and the composites confirmed that -SiC forms as matrix phase. Pyrolysis temperature plays a role on the crystallite size and the mechanical properties of the composites. Examination of the interior of the composite and cross-section clearly shows the matrix formation around individual fibers in the mini-composite. Mechanical property evaluation showed that the tensile strength of the composites fabricated at 1600 ◦ C gave substantially higher values and showed typical composite failure behavior compared to those fabricated at 1700 ◦ C. A detailed investigation involving Weibull statistical analysis revealed that the composites fabricated at 1600 ◦ C, gave not only higher tensile strength and fracture energy, but also showed higher Weibull modulus indicative of less scatter in the data. This may be due to the micro-crystalline nature of the matrix formed at 1600 ◦ C.
Acknowledgements The authors would like to acknowledge the help from the XRD group for carrying out XRD. They are thankful to Director, Defence Metallurgical Research Laboratory for his constant encouragement. They also acknowledge the finanacial support from Defence Research and Development Organisation, Government of India to carry out the present study.
references
Boisvert, R.P., Diefendorf, R.J., 1988. Polymeric precursor SiC matrix composites. Ceram. Eng. Sci. Proc. 9 (7–8), 873–880. Cheng, L., Xu, Y., Zheng, L., Yin, X., 2001. Effect of carbon interlayer on oxidation behavior of C/SiC composites with a coating from room temperature to 1500 ◦ C. Mater. Sci. Eng. A 300, 219–225. Clegg, W.J., 2000. The role of carbon in the sintering of boron-doped SiC. J. Am. Ceram. Soc 8 (5.), 1039–1043. Davidge, R.W., 1979. Mechanical Behavior of Ceramics. Cambridge University press, Cambridge, pp. 131–140. Evans, A.G., 1989. The machanical properties of fiber reinforced caremic matrix composites. Mater. Sci. Eng. A 158, 227–239. Fitzer, E., Gadow, R., 1986. Fabrication of carbon/siliconcarbide composites by isothermal chemical vapor infiltration, using the in-situ whisker growing and matrix filling process. Am. Ceram. Soc. Bull. 65, 326. Gonczy, S.T., Sprandel, R.C., Faber, K.T., Robert, R., 1997. Tensile tests of miniature fiber reinforced ceramic composites for screening process–property relations. Ceram. Eng. Sci. Proc. 18 (3), 729–736. He, X.-B., Zhang, X.-M., Zhang, C.-R., Zhou, X.-G., Zhou, A.-C., 2001. Micro structures of a carbon-fiber-reinforced silicon-carbide composites produced by precursor pyrolysis and hot pressing. Combust. Sci. Technol. 61, 117–123. Kenji, O., Yoshitsugu, K., Tsugio, I., Koichi, N., 2001. Carbon fiber reinforced silicon nitride based nanocomposite materials and
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method for preparing same. European Patent No. 053,686, 6A1. Kim, Y.-W., Song, J.-S., Park, S.-W., Lee, J.-G., 1993. Nicalon-fibre-reinforced silicon-carbide composites via polymer solution infiltration and chemical vapour infiltration. J. Mater. Sci. 28, 3866–3868. Krenkel, W., 2001. Cost effective processing of CMC composites by melt infiltration (LSI-process). Ceram. Eng. Sci. Proc. 22 (3), 443–470. Lamon, J., Lissart, N., Rechiniac, C., 1993. Micromechanical and statistical approach to the behaviour of CMC’s. Ceram. Eng. Sci. Proc. 14 (9–10), 115–1124. Lamon, J., Rebillat, F., Evans, A.G., 1995. Assessment of a micro composite test procedure for evaluating constituent properties of ceramic matrix composites. J. Am. Ceram. Soc. 78 (2), 401. Larsen, David C., Stuchly, Sharon L., 1990. The mechanical evaluation of ceramic fiber composites. In: Mazdiyasni, K.S. (Ed.), Fiber reinforced ceramic composites, Materials, Processing and Technology. Noyes Publications, USA, pp. 182–221. Mentz, Juliane, Muller, Marcus, 2006. Carbon-fiber reinforced carbon composites filled with SiC particles forming a porous matrix. Mater. Sci. Eng. A 64–69, 425. Naslain, Roger, 2004. Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview. Compos. Sci. Technol. 64, 155– 170. Naslain, R., Lamon, J., Pailler, R., Bourrat, X., Guette, A., Langlais, F., 1999. Micro/mini composites: a useful approach to the design and development of non-oxide CMCs. Compos.: Part A 30, 537–547. Naslain, R., Pailler, R., Bourrat, X., Bertarnd, S., Lamouroux, F., 2001. Synthesis of highly tailored ceramic matrix composites by pressure-pulsed CVI. Solid State Ionics 541–548, 141–142. Prewo, K.M., 1989. Fiber reinforcedceramics-new oppourtunities for composite materials. Am. Ceram. Bull. 68 (1–4), 395. Prewo, K., Brennan, J.J., 1980. High strength silicon carbide fiber-reinforced glass-matrix composites. J. Mater. Sci. 15, 463–485. Sato, K., Tezuka, A., Funayama, O., Isoda, T., Terada, Y., Kato, S., Iwata, M., 1999. Fabrication and pressure testing of a gas-turbine component manufacturred by a pre ceramic–polymer-impregnation method. Combust. Sci. Technol. 59, 853–859. Shetty, D.K., Pascucci, M.R., Mutsuddy, B.C., Wills, R.R., 1985. SiC mono-filament composites reinforced silicon nitride matrix composites. Ceram. Eng. Sci. Proc. 6 (78), 632–645. Stinton, D.P., Besmann, M.T., Lowdon, R.A., 1985. Advanced ceramics by chemical vapour deposition technique. Ceram. Bull. 67, 350–355. Stinton, D.P., Caputo, A.J., Lowden, R.A., 1986. Synthesis of fiber-reinforced SiC composites by chemical vapor infilrtaion. Ame Ceram. Soc. Bull. 65 (2), 347–350. Subrahmanyam, J., Mohan Rao, R., Ghosal, P., Vijay Kumar, M., 2004. A process for the preparation of reinforced ceramic matrix composites having both the advantages of high strength and high fracture toughness, Indian Patent applied. 406/DEL/2004. Tressler, R.E., 1986. Tailoring multiphase and composite ceramics. Mater. Sci. Res. 20, 145–164. Weimer, A.W., Nilsen, K.J., Cochran, G.A., Roach, R.P., 1993. Kinetics of carbothermal reduction synthesis of beta silicon carbide. AlChE J. 39 (3), 493–503. Zhong, G., Sano, H., Uchiyama, Y., Chang, H., 1998. Preparation and fracture behavior of carbon fiber/SiC composites by multiple impregnation and pyrolysis of polycarbosilane. J. Ceram. Soc. Jpn. 106 (2), 1155–1161.
journal homepage: www.elsevier.com/locate/jmatprotec
Carbon fiber reinforced silicon carbide mini-composites-solution approach N. Padmavathi a,b,∗ , J. Subrahmanyam a , K.K. Ray b , R. Mohanrao a , P. Ghosal a , Sweety Kumari a a b
Defence Metallurgical Research Laboratory, Kanchanbagh and Hyderabad 58, India Metallurgical and Materials Science Engineering, Indian Institution of Technology, Kharagpur 721302, India
a r t i c l e
i n f o
a b s t r a c t
Article history:
Carbon fiber reinforced composites are widely used as aerospace materials owing to their
Received 4 July 2007
superior properties like high specific strength and high fracture toughness. Oxidation above
Received in revised form
600 ◦ C is one of the major problems encountered with carbon–carbon composites. Replace-
1 November 2007
ment of carbon matrix with SiC is one of the methods to improve the oxidation resistance.
Accepted 12 November 2007
A relatively inexpensive process is developed for the fabrication of carbon fiber reinforced SiC composites through solution approach. This process allowed forming high purity and fine-grained matrix in the fiber reinforced composites. Carbon fiber tow mini-composites
Keywords:
were fabricated using this process. Typical tensile strength and fracture energy of these com-
Carbon fiber
posites were 316 MPa and 0.74 MJ/m3 at room temperature (RT). The composites exhibited
SiC
load-carrying capability after crack initiation.
Solution approach
© 2007 Elsevier B.V. All rights reserved.
Composite Tensile strength
1.
Introduction
Fiber-reinforced-ceramic-matrix composites have recently attracted attention for use in high-temperature structural applications (Stinton et al., 1986; Prewo, 1989). The primary reason for this interest lies in the assumption that strong ceramic fibers can prevent catastrophic brittle failure in ceramics by providing various energy-dissipation processes during crack advance (Evans, 1989). Early works utilizing ceramic fibers in ceramic matrices demonstrated the potential of this approach (Prewo and Brennan, 1980; Stinton et al., 1985, 1986; Shetty et al., 1985; Naslain et al., 2001; Tressler, 1986; Fitzer and Gadow, 1986). Carbon fiber reinforced silicon carbide (C/SiC) composites provide excellent thermo-mechanical properties at tempera-
tures up to 2000 ◦ C (Zhong et al., 1998; He et al., 2001; Mentz and Muller, 2006). Their high strength-to-weight ratio, which sustains at the high temperatures, makes them potential candidates for highly demanding engineering applications such as heat shields and structural components for re-entry space vehicles, high performance brake discs, and ultra-high temperature heat exchanger tubes. For these elements, C/SiC composites have the most favourable properties and are, therefore, at the challenging front of the contemporary materials research. However, a generic problem that must be overcome is that ceramic fabrication processes tend to mechanically and chemically damage the fibers when they are consolidated within a ceramic matrix. For example, the fibers may be broken by a pressing operation, or the high sintering temperature required
∗ Corresponding author at: Defence Metallurgical and Research Laboratory, Kanchanbagh and Hyderabad 58, India. Tel.: +91 9440413314; fax: +91 40 24340683. E-mail address: nandigam [email protected] (N. Padmavathi). 0924-0136/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.11.124
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
to densify the ceramic matrix may damage the fibers or cause them to react chemically with the matrix. C/SiC composites can be fabricated by a CVI infiltration process, yet their high performance is attained at prohibiting costs for many applications (Kim et al., 1993; Kenji et al., 2001; Cheng et al., 2001). Polymer impregnation and pyrolysis (PIP) derived SiC matrix contains free carbon and oxygen (Sato et al., 1999; Boisvert and Diefendorf, 1988). Liquid silicon infiltration processes also a common fabrication route, offer some cost advantage, but produces SiC matrix along with free silicon (Krenkel, 2001). New processing techniques are to be developed to get high quality SiC matrix at low cost. Soft solution processing described in this paper allows fabrication of advanced fiber reinforced composites using aqueous solutions, without expensive equipment, providing an environmental friendly route. In this study, carbon fiber reinforced SiC composites have been prepared by a solution approach (Weimer et al., 1993). Solution approach or soft chemistry route involves the low temperature modification of existing solid structures to form new solids that retains many of the structural features of the precursor phase. A mixture of colloidal silica and sucrose in solution form were used as a precursor for formation of SiC matrix. Drying and heating to 500 ◦ C converts the sucrose into carbon and SiC forms by the carbothermic reduction of silica at high temperatures by the following reaction, SiO2 + 3C → SiC + 2CO
(1)
Mini-composites were fabricated by impregnating the solution phase into a carbon fiber tow, drying, converting sucrose into carbon phase and finally reducing silica with carbon to form SiC. The principal advantage of this approach is the relatively inexpensive ingredients that constitute the solution
435
phase, to get fine-grained SiC matrix in a fiber reinforced composite (Subrahmanyam et al., 2004).
2.
Experimental procedure
2.1.
Raw materials
The carbon fiber used in the study was supplied by Tae Kwang Ind. Co. Ltd and Japan, Grade-TZ-300 with 12,000 filaments and filament diameter of 6 m (12 K). Colloidal silica was obtained from Bee Chem. Chemicals Company, Kanpur and India, and Sucrose from Qualigens Fine chemicals, India.
2.2.
Composite-synthesis
Grade-TZ-300 carbon fiber tow was initially washed with acetone to remove impurities and sizing. The composition of the precursor solution is 0.25 M of colloidal silica and 0.876 M of sucrose in distilled water. The fiber tows were vacuum impregnated with precursor solutions of SiC (20–25 ml per 80 mg of fiber tow) to obtain mini-composites with SiC matrix. Impregnated tows were dried carefully at room temperature and at 60–70 ◦ C for 12 h. Dried tows were carbonized at 500 ◦ C for 1 h for converting sucrose into carbon. These steps were repeated for five to six cycles. After the desired impregnation/heat treatment cycles, carbonized tows were pyrolysed at 1600 and 1700 ◦ C for 3 h under flowing argon to obtain the SiC matrix. Flow sheet for the process is shown in Fig. 1. Bulk densities of the samples were measured with deionised water as immersion medium according to the Archimedes principle. The densities of the composites increased with the number of infiltration cycles. Further cycles were not recommended after completing of six cycles due to
Fig. 1 – Flow sheet for the process.
436
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
the rapid decreasing effectiveness of further infiltration. Density of the composite was 1940 kg cm−3 and porosity is 30%. The volume fraction of fiber and matrix are 88% and 12%, respectively.
2.3.
Microstructural characterization
Philips Model No. PW 320 is used to carry out X-ray diffraction (XRD) of samples to identify the crystalline phases of the powder samples produced from the precursor solution as well as the phases from mini-composites. The microstructures of the mini-composites were examined using an optical microscope as well as a scanning electron microscopy (Model No. 7060, SEM- Oxford Leo 440i) equipped with an energy dispersive X-ray spectrometer (EDS).
2.4.
Evaluation of mechanical properties
Relatively simple bend testing (Flexure testing) is generally used for ceramic samples. However it is not suitable for ceramic matrix composites due to non-uniform stress state and brittle nature of composites matrix. In the present work mini-composite approach, single or a few tows are combined together and impregnated to get the composite (Naslain et al., 1999; Lamon et al., 1995). The benefits with the minicomposite approach are better utilization of the expensive fiber material and the mini-tow composite tensile sample has a uni-axial tensile stress condition, which gives the more appropriate data, for process development compared to flexure testing and it allows study of various process variables in a shorter time (Gonczy et al., 1997; Larsen and Stuchly, 1990). The mechanical tests were performed using a servohydraulic testing machine (Model 8801; Instron Corp., UK) at room temperature under a displacement rate of 0.5 mm/min. The load was measured using a 1000 N load cell. Specimens are prepared by following the procedure outlined for minicomposites at room temperature with a gauge length of 50 mm and using paper tabs for gripping with a suitable adhesive (Naslain et al., 1999). Once the mini-composite sample was mounted on the machine, the cardboard tab supporting the fiber was cut and testing carried out. The properties of mini-composites are characterized by two separate mechanical parameters in this study: (1) tensile strength was calculated from the stress–strain curve of each composite and (2) fracture energy was measured by area under the curve for each composite. Fracture energy gives the damage tolerance of the composite. Weibull modulus (m) reflects the degree of variability in the strength values; where a higher number indicates less scatter in data. For ceramics, the values are in the range of 5–20 (Davidge, 1979; Lamon et al., 1993).
3.
Results and discussion
3.1.
Evaluation of the microstructures
XRD pattern of carbon fiber tow for comparison with tow composite is presented in Figs. 2 and 3 shows that the XRD pattern of the powder samples produced from the precursor
Fig. 2 – XRD pattern for carbon fiber tow.
solution, which confirms the formation of -SiC. However, the powders prepared at 1600 and 1700 ◦ C not showing any difference in terms of peak intensity. Fig. 4 shows the XRD patterns of the mini-composites, synthesized at two different pyrolysis temperatures. The C-SiC composite consists mainly -SiC and minor graphite peaks originating from the fiber. There is no trace of unreacted SiO2 in diffraction patterns. From the diffractogram (Fig. 4), it is clear that at a pyrolysis temperature of 1600 ◦ C, -SiC peaks are just appearing and at 1700 ◦ C, the peak intensity increased considerably. A graphite peak is also present in these diffractograms indicating the presence of carbon fibers. This may be due to the fact that at 1600 ◦ C, micro-crystalline -SiC forms from an amorphous solution phase and at 1700 ◦ C the size of the crystallites increases (Naslain, 2004).
Fig. 3 – XRD pattern for the powders prepared through soft solution route (a) at 1600 ◦ C and (b) at 1700 ◦ C/3 h.
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Fig. 4 – XRD pattern for the C/SiC composites prepared through soft solution route pyrolysed at (a) 1600 ◦ C and (b). 1700 ◦ C for duration of 3 h.
3.2.
Microstructures of composites
Fig. 5 shows the SEM pictures of carbon fiber tow and mini-composites which is opened vertically to view the microstructure of the interior surface of the composite. It shows the distribution of matrix between the fibers through the composite. It is difficult to prepare cross-sectional metallographic samples from the mini-composites. During the preparation of the porous and brittle matrix of mini-composite
Fig. 6 – Cross-sectional electron micrographs of composite: (a) SEM micrograph and (b) optical micrograph with polarized light.
was severely damaged. During dry cutting using a diamond saw without coolant, the brittle fibers were torn and pulled out. When cut in wet condition, the removed carbon residue and fiber remnants were smeared on the surface, making further observations difficult. Finally it was possible to view the microstructure of the composites by filling the porosity with a low-viscosity resin. This fixed the individual fibers and facilitated polishing of the sample. And the charged surfaces of resin portions were carefully avoided while observing the polished section of the mini-composites. The cross section of the polished composite was observed under SEM. Fig. 6(a) shows a typical micrograph of pyrolysed CMC showing distribution of fibers and matrix. Fig. 6(b) shows the optical micrograph using polarized light of minicomposite showing individual fibers clearly separated by the SiC matrix.
3.3.
Fig. 5 – SEM micrograph of (a) carbon fiber and (b) C/SiC composites fabricated by solution approach.
Mechanical properties
Tensile tests were performed on carbon fiber tows and the mini-composites. Fig. 7(a) shows the stress–strain curve for carbon fiber tow at room temperature. The fiber tow showed gradual decrease in load after the point of maximum load, due to individual fiber breakage. Fig. 7(b) shows the tensile behavior of two composites, one pyrolysed at 1600 ◦ C and the other processed at 1700 ◦ C. The composites pyrolysed at 1600 ◦ C showed a gradual decrease in stress after the point
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j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 4 ( 2 0 0 8 ) 434–439
Fig. 7 – Stress–strain curves at room temperature (a) carbon fiber tow and in (b). (1) Pyrolysed at 1600 ◦ C and (2) At 1700 ◦ C.
Fig. 8 – Plot to obtain the Weibull modulus of 6.6 for 14 samples pyrolysed at 1600 ◦ C tested at room temperature. Fig. 9 – Plot to obtain the Weibull modulus of 3.7 for 15 samples pyrolysed at 1700 ◦ C tested at room temperature. of maximum load, similar to the fiber tow, typical of a composite failure. The mini-composites pyrolysed at 1700 ◦ C, not showing this type of behavior. Stress decreased rapidly after the point of maximum load indicative of catastrophic monolithic ceramic type failure behavior. A possible reason for this difference could be due to the difference between the matrices obtained at these two different temperatures. As shown in the XRD patterns for these composites, the SiC matrix obtained at 1600 ◦ C is in a micro-crystalline form with low Xray peak intensities. On the other hand the peak intensities have increased considerably for the SiC matrix processed at 1700 ◦ C showing the evidence for larger crystallite size. In order to quantify these results, several mini-composite samples were prepared at both temperatures and tensile tests were carried out. Weibull statistical analysis was carried out on this data to obtain the Weibull modulus. Weibull modulus gives an idea about the scatter in the data, higher the value lowers the scatter. Figs. 8 and 9 present the Weibull
plots for the composites processed at 1600 and 1700 ◦ C, respectively. Fracture energy for each of these composites is obtained from the area under the tensile curves. The average values of the fracture energy, along with the tensile strength and the Weibull modulus are presented in Table 1. As can be seen from Table 1, carbon fiber tows without any matrix shows highest strength, fracture energy and Weibull modulus. It is well known that the fiber strength cannot be realized in the composites due to (i) fiber damage during composite fabrication process and (ii) constraint imposed by the matrix on the fiber. Among the composites, the one processed at 1600 ◦ C, gave not only substantially higher tensile strength and fracture energy, but also higher Weibull modulus indicative of less scatter in the data. As already mentioned, this may be due to the micro-crystalline nature of the matrix formed at 1600 ◦ C. In addition silicon carbide can react with the residual
Table 1 – Mechanical properties of C/SiC composites Material Carbon C-SiC at 1600 ◦ C for 3 h C-SiC at 1700 ◦ C for 3 h
Tensile strength (MPa) 1146 ± 104 316 ± 48 149 ± 38
Fracture energy (MJ/m3 ) 10.4 0.74 0.07
Weibull modulus (m) 10.5 6.6 3.7
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silica in the matrix at high temperature resulting in weaker matrix (Clegg, 2000).
4.
Conclusions
This investigation has clearly demonstrated the possibility of fabrication of the carbon fiber reinforced SiC mini-composites by soft solution approach. X ray diffraction of the powders and the composites confirmed that -SiC forms as matrix phase. Pyrolysis temperature plays a role on the crystallite size and the mechanical properties of the composites. Examination of the interior of the composite and cross-section clearly shows the matrix formation around individual fibers in the mini-composite. Mechanical property evaluation showed that the tensile strength of the composites fabricated at 1600 ◦ C gave substantially higher values and showed typical composite failure behavior compared to those fabricated at 1700 ◦ C. A detailed investigation involving Weibull statistical analysis revealed that the composites fabricated at 1600 ◦ C, gave not only higher tensile strength and fracture energy, but also showed higher Weibull modulus indicative of less scatter in the data. This may be due to the micro-crystalline nature of the matrix formed at 1600 ◦ C.
Acknowledgements The authors would like to acknowledge the help from the XRD group for carrying out XRD. They are thankful to Director, Defence Metallurgical Research Laboratory for his constant encouragement. They also acknowledge the finanacial support from Defence Research and Development Organisation, Government of India to carry out the present study.
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