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Fabrication and microstructural characterization of porous silicon carbide with nano-sized powders
Materials Science and Engineering B 148 (2008) 211–214
Fabrication and microstructural characterization of porous silicon carbide with nano-sized powders Manabu Fukushima ∗ , You Zhou, Yu-Ichi Yoshizawa National Institute of Advanced Industrial Science and Technology (AIST), 2266-98 Shimo-Shidami, Moriyama-ku, Nagoya 463-8560, Japan Received 22 May 2007; received in revised form 31 July 2007; accepted 3 September 2007
Abstract Porous silicon carbide was fabricated by using nano-sized SiC powder additions and different cold isostatic pressing (CIP) conditions followed by sintering at 1500–1800 ◦ C. The relationship between the processing conditions, pore size and microstructure was examined. The cold isostatic pressing conditions, sintering temperature and nano-sized additives were effective for controlling pore size and microstructure. The pore size and particle size increased with increasing sintering temperature, attributed to surface diffusion. However, no densification occurred because of pore enlargement. In addition, the compressive strength increased with sintering temperature and reached values as high as 513 MPa. This was due to the formation of well-developed neck areas. This study suggests that the promoted mass transfer can provide high strength due to increased neck area. © 2007 Elsevier B.V. All rights reserved. Keywords: Silicon carbide; Porous ceramic; Pore size distribution; Surface diffusion; Filter and membrane support
1. Introduction Macroporous silicon carbide (SiC) ceramics have played an important role in many fields of energy production and environmental protection such as hydrogen permselective membrane supports, diesel particulate filters, slurry reuse after polishing and water purification, because of their low thermal-expansion coefficient and good thermal-shock resistance as well as excellent mechanical and chemical stability at elevated temperatures [1–6]. For these filter applications, high porosity and precise pore size control are required. In addition, strength is an important factor for filtration efficiency, because high strength can allow high-pressure operation. The strength of a porous ceramic is found to be largely dependent on the neck thickness [7]. Recently, we have successfully developed porous SiC membrane supports with and without alumina, which had pore sizes of around 0.1–0.4 m [6]. For these porous SiC with alumina additives, the alumina reacted with the silica layer existing naturally on the surface of the raw SiC particles to form an Al2 O3 –SiO2 liquid phase, which covered the SiC particles and retarded mass transfer and grain growth, thereby retaining the
∗
Corresponding author. Tel.: +81 52 736 7161; fax: +81 52 736 7405. E-mail address: [email protected] (M. Fukushima).
0921-5107/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2007.09.026
small pore size of the green body. In contrast, in the microstructure of porous SiC without additives, the neck area and particle size were enlarged. When these phenomena occur, the fine particles contained in the starting powder become important factors in controlling the microstructure, since fine particles have large surface energy and promote mass transfer [6,8]. Thus, previous studies suggest that the neck area can be enlarged by the existence of fine particles. In this paper, porous SiC with nano-sized powder additives were prepared, and the relationship between the processing conditions, properties and microstructure was investigated using scanning electron microscopy (SEM) and pore size distributions. 2. Experimental procedure High purity -SiC (Ibiden Co. Ltd., Gifu, Japan) powder with an average particle size of 0.30 m and a specific surface area of 20.0 m2 /g was used as the raw material. The main impurities of the SiC powder were 0.25 mass% oxygen and 1 mass% free carbon. As the additive, high purity -SiC (Sumitomo Osaka Cement Co. Ltd., Tokyo, Japan) with an average particle size of 30 nm and a specific surface area of 40–50 m2 /g was used. The main impurities of the nano-sized powder were 0.3–0.6 mass% oxygen and 1.8–2.4 mass% free carbon. Powder mixtures with weight ratios of nano-sized SiC/micro-sized SiC = 10/90 were
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blended in ethanol for 1 h by using a planetary mill with a SiC pot and SiC balls. After mixing, the slurry was dried with a rotary evaporator and a vacuum oven at 110 ◦ C for 4 h, and screened through a 110 mesh sieve. The powder compacts were formed to a disk shape without any binder by using a steel mold at 40 MPa, and they were then subjected to cold isostatic pressing (CIP) at either 100 or 400 MPa (hereafter identified as CIP100 or CIP400, respectively). The size of the formed disk specimens was approximately 20 mm in diameter and 5 mm in thickness. The green bodies were placed in a capped graphite crucible, and sintered at 1500, 1600, 1700 and 1800 ◦ C for 2 h under Ar gas flow. The heating rate from room temperature to 1200 ◦ C was 20 ◦ C/min, and 15 ◦ C/min from 1200 ◦ C to setting temperatures. The apparent densities of the sintered specimens were measured by water displacement using the Archimedes method, according to the equation [9]: d=
ρWa , Ww − W l
where Wl , Ww and Wa indicate the weight in water, wet weight and dry weight, respectively, and ρ is the water density at the measurement temperature. Green body densities were calculated from the dimensions and weight of the specimens. Specific surface area was measured by nitrogen adsorption/desorption (Yuasa Ionics Inc., Autosorb, Osaka, Japan), where surface area was determined from a BET (Brunauer, Emmet and Teller) analysis in the P/P0 range of 0.05–0.30 using a molecular cross-sectional area for N2 of 0.163 nm2 and 10 points [10]. The microstructures were observed using a scanning electron microscopy (SEM; JEOL-6330F, Japan). The pore size distributions were measured by mercury porosimetry (Yuasa Ionics Inc., PoreMaster-GT, Osaka, Japan) in the range of 6.1 nm–426 m. The weight of the specimens used for the measurement was around 0.5 g. The compressive strength was measured according to JIS R1608. The used cylindrical specimen had a dimension of 5 mm and 12.5 mm in height and was fractured using a testing machine (MTS Systems Corporation, Sintech 10/GL, Minnesota, USA). The load was applied on the two parallel surfaces at a cross-head speed of 0.5 mm/min. For every material, five specimens were tested.
Table 1 Linear shrinkage and relative density vs. sintering temperature CIP pressure
Sintering temperature (◦ C)
Linear shrinkage (%)
Relative density (%)
100 MPa
Green body 1500 1600 1700 1800
– 0.5 0.5 0.5 0.5
55 56 56 56 56
400 MPa
Green body 1500 1600 1700 1800
– 0.2 0.2 0.2 0.2
60 61 61 61 61
Fig. 1 shows typical examples of the pore size distributions of the specimens treated with different CIP pressures and sintered at different temperatures. The pore sizes of the specimens sintered at 1500 ◦ C were in the ranges of 0.10–0.22 and 0.09–0.19 m for CIP100 and CIP400 specimens, respectively. When sintered at 1800 ◦ C, the pore sizes increased to 0.50–1.15 m for CIP100 and 0.40–0.94 m for CIP400, respectively. Hence, higher CIP pressure reduced the pore size and high sintering temperature resulted in enlarged pore sizes. The peak pore size shown in Fig. 2 demonstrates a clear relationship between pore size, sintering temperature, nano-sized powder addition, and CIP pressure. The peak sizes for CIP100 were 0.19, 0.50, 0.75 and 0.95 m from 1500 to 1800 ◦ C, and those for CIP400 were 0.14, 0.41, 0.57 and 0.73 m, respectively. Table 2 shows the typical values of BET surface area and compressive strength of the specimens treated at CIP400 and sintered at 1500 and 1800 ◦ C. The specimens sintered at 1500 ◦ C had a surface area of 9.3 m2 /g. However, when sintered at 1800 ◦ C, the surface area decreased to 1.9 m2 /g, which was about 20% of that of the specimen sintered at 1500 ◦ C. Since the porosity measured by the Archimedes method was constant during sintering, the decrease in surface area is considered to be due to the decrease of the number of fine particles, grain growth
3. Results and discussion The linear shrinkage and relative density as a function of sintering temperature are listed in Table 1. The specimens did not exhibit significant linear shrinkage, regardless of sintering temperature. An effect of CIP pressure on density was clearly observed; higher pressure resulting in higher density. For all sintering temperatures, the specimens treated at CIP400 MPa exhibited 5% higher relative density than those treated at CIP100 MPa. In addition, densities were constant regardless of sintering temperature. These observed results are consistent with previous reports of -SiC compacts produced without any sintering aids [6,11]. Hase et al. have reported that heating a compact of -SiC with an average particle size of 0.1 m to 2100 ◦ C resulted in no densification or shrinkage.
Fig. 1. Pore size distributions of porous SiC sintered at 1500 and 1800 ◦ C.
M. Fukushima et al. / Materials Science and Engineering B 148 (2008) 211–214
Fig. 2. Peak pore size of porous SiC sintered at various temperatures. Table 2 Typical values of BET surface area and compressive strength of the specimens treated at CIP400 and sintered at 1500 and 1800 ◦ C Sintering temperature (◦ C)
BET surface area (m2 /g)
Compressive strength (MPa)
1500 1800
9.3 ± 0.6 1.9 ± 0.2
122 ± 22 513 ± 38
Fig. 3. SEM micrographs of the fractured surfaces of the CIP400 specimens sintered at (a) 1500 and (b) 1800 ◦ C, respectively.
213
and enlarged pore size. The results of strength tests showed that the compressive strength of the specimens sintered at 1500 ◦ C was 122 MPa, and this increased to 522 MPa when sintering at 1800 ◦ C. Although the pore size of the specimen sintered at 1800 ◦ C was 5 times larger than that sintered at 1500 ◦ C, the increase of sintering temperature resulted in increased strength, suggesting that well-developed neck areas were formed at higher sintering temperatures. Fig. 3 shows SEM micrographs of the fractured surfaces of the CIP400 specimens sintered at (a) 1500 and (b) 1800 ◦ C, respectively. It is clearly seen that the microstructure changed with sintering temperature. Fine particles and small pores were observed in the specimens sintered at 1500 ◦ C but disappeared after sintering at 1800 ◦ C, where appreciable grain growth and enlarged pores were observed. These phenomena can be more clearly seen in high magnification observations, as shown in Fig. 4. In addition, large necks were observed in the specimen sintered at 1800 ◦ C. In general, the surface of the SiC particles is covered with a thin, naturally existing, SiO2 layer. When the green compact is heated above 1400 ◦ C, this surface SiO2 layer can react with free carbon (one of the main impurities) or SiC, resulting in
Fig. 4. SEM micrographs of the fractured surfaces of the CIP400 specimens sintered at (a) 1500 and (b) 1800 ◦ C, respectively.
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the evolution of SiO and CO gas, according to the following equations [4,11–14]: SiO2 + C → SiO(g) + CO(g) 2SiO2 + SiC → 3SiO(g) + CO(g) SiO(g) + 2C(s) → SiC(s) + CO(g) Through these reactions a silica free surface is formed on the SiC particles, which can accelerate the surface diffusion. The diffusion process from the particle surface to the neck regions should enlarge the neck area. This may result in the consumption of the nano-sized particles, since they have larger surface energy and smaller diffusion distance than submicron-sized particles. Mass transfer from the surface towards the neck should result in an increase in both pore size and grain size as well as the neck area. Compared to our previous reports, the neck area observed in this study was well developed, which was due to the promoted mass transfer by nano-sized powder additions [15]. The present compressive strength (513 MPa) was also much higher than the strength reported previously (15 MPa) for our previous porous SiC with similar density (58%) and peak pore size (1.7 m). Thus, this study suggests that fabricating porous SiC by promoting mass transfer can provide high strength due to increased neck area. 4. Conclusion Porous silicon carbide sintered at 1500–1800 ◦ C was prepared by using different CIP pressures and nano-sized SiC additives. The role of these processing factors on pore size and microstructure was investigated. A higher CIP pressure treatment was effective in reducing pore size. The pore size and particle size increased with increasing sintering temperature, and
the disappearance of the added nano-sized powder was observed. Enlarged neck regions were also found, due to mass transfer of the nano-sized particles from the particle surface toward the neck area. The compressive strength increased with sintering temperature and attained a value over 500 MPa through the formation of well-developed neck regions. The addition of nano-sized powder suggests that the promoted mass transfer from the surface to neck regions can provide high strength due to increased neck area. References [1] M.J. Ledoux, S. Hantzer, C.P. Huu, J. Guille, M.P. Desaneaux, J. Catal. 114 (1988) 176–185. [2] M. Benaissa, J. Werckmann, G. Ehret, E. Peschiera, J. Guille, M.J. Ledoux, J. Mater. Sci. 29 (1994) 4700–4707. [3] N. Keller, C. Pham-Huu, S. Roy, M.J. Ledoux, C. Estournes, J. Guille, J. Mater. Sci. 34 (1999) 3189–3202. [4] V. Suwanmethanond, E.P. Goo, K.T. Liu, G. Johnston, M. Sahimi, T.T. Tsotsis, Ind. Eng. Chem. Res. 39 (2000) 3264–3271. [5] P.-K. Lin, D.-S. Tsai, J. Am. Ceram. Soc. 80 (1997) 365–372. [6] M. Fukushima, Y. Zhou, Y. Iwamoto, S. Yamazaki, T. Nagano, H. Miyazaki, Y. Yoshizawa, K. Hirao, J. Am. Ceram. Soc. 89 (2006) 1523–1529. [7] J. She, J.-F. Yang, N. Kondo, T. Ohji, S. Kanzaki, Z.-Y. Deng, J. Am. Ceram. Soc. 85 (2002) 2852–2854. [8] F. Wakai, M. Yoshida, Y. Shinoda, T. Akatsu, Acta Mater. 53 (2005) 1361–1371. [9] J.T. Jones, M.F. Berard, Ceramics—Industrial Processing and Testing, The Iowa State University Press, Ames, Iowa, 1993, pp. 172–174. [10] S. Brunauer, P.H. Emmett, E. Teller, J. Am. Chem. Soc. 60 (1938) 309–319. [11] T. Hase, H. Suzuki, I. Tomozuka, Yogyo-Kyoukai-shi 87 (1979) 317–321 (in Japanese). [12] C. Greskovich, J.H. Rosolowski, J. Am. Ceram. Soc. 59 (1976) 336–343. [13] T. Hase, H. Suzuki, Yogyo-Kyoukai-shi 88 (1980) 258–264 [in Japanese]. [14] T. Grande, H. Sommerset, E. Hagen, K. Wiik, M.-A. Einarsrud, J. Am. Ceram. Soc. 80 (1997) 1047–1052. [15] M. Fukushima, Y. Zhou, Y. Yoshizawa, H. Miyazaki, K. Hirao, Presented at the 31st International Cocoa Beach Conference & Exposition on Advanced Ceramics & Composites, 2007.
Fabrication and microstructural characterization of porous silicon carbide with nano-sized powders Manabu Fukushima ∗ , You Zhou, Yu-Ichi Yoshizawa National Institute of Advanced Industrial Science and Technology (AIST), 2266-98 Shimo-Shidami, Moriyama-ku, Nagoya 463-8560, Japan Received 22 May 2007; received in revised form 31 July 2007; accepted 3 September 2007
Abstract Porous silicon carbide was fabricated by using nano-sized SiC powder additions and different cold isostatic pressing (CIP) conditions followed by sintering at 1500–1800 ◦ C. The relationship between the processing conditions, pore size and microstructure was examined. The cold isostatic pressing conditions, sintering temperature and nano-sized additives were effective for controlling pore size and microstructure. The pore size and particle size increased with increasing sintering temperature, attributed to surface diffusion. However, no densification occurred because of pore enlargement. In addition, the compressive strength increased with sintering temperature and reached values as high as 513 MPa. This was due to the formation of well-developed neck areas. This study suggests that the promoted mass transfer can provide high strength due to increased neck area. © 2007 Elsevier B.V. All rights reserved. Keywords: Silicon carbide; Porous ceramic; Pore size distribution; Surface diffusion; Filter and membrane support
1. Introduction Macroporous silicon carbide (SiC) ceramics have played an important role in many fields of energy production and environmental protection such as hydrogen permselective membrane supports, diesel particulate filters, slurry reuse after polishing and water purification, because of their low thermal-expansion coefficient and good thermal-shock resistance as well as excellent mechanical and chemical stability at elevated temperatures [1–6]. For these filter applications, high porosity and precise pore size control are required. In addition, strength is an important factor for filtration efficiency, because high strength can allow high-pressure operation. The strength of a porous ceramic is found to be largely dependent on the neck thickness [7]. Recently, we have successfully developed porous SiC membrane supports with and without alumina, which had pore sizes of around 0.1–0.4 m [6]. For these porous SiC with alumina additives, the alumina reacted with the silica layer existing naturally on the surface of the raw SiC particles to form an Al2 O3 –SiO2 liquid phase, which covered the SiC particles and retarded mass transfer and grain growth, thereby retaining the
∗
Corresponding author. Tel.: +81 52 736 7161; fax: +81 52 736 7405. E-mail address: [email protected] (M. Fukushima).
0921-5107/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2007.09.026
small pore size of the green body. In contrast, in the microstructure of porous SiC without additives, the neck area and particle size were enlarged. When these phenomena occur, the fine particles contained in the starting powder become important factors in controlling the microstructure, since fine particles have large surface energy and promote mass transfer [6,8]. Thus, previous studies suggest that the neck area can be enlarged by the existence of fine particles. In this paper, porous SiC with nano-sized powder additives were prepared, and the relationship between the processing conditions, properties and microstructure was investigated using scanning electron microscopy (SEM) and pore size distributions. 2. Experimental procedure High purity -SiC (Ibiden Co. Ltd., Gifu, Japan) powder with an average particle size of 0.30 m and a specific surface area of 20.0 m2 /g was used as the raw material. The main impurities of the SiC powder were 0.25 mass% oxygen and 1 mass% free carbon. As the additive, high purity -SiC (Sumitomo Osaka Cement Co. Ltd., Tokyo, Japan) with an average particle size of 30 nm and a specific surface area of 40–50 m2 /g was used. The main impurities of the nano-sized powder were 0.3–0.6 mass% oxygen and 1.8–2.4 mass% free carbon. Powder mixtures with weight ratios of nano-sized SiC/micro-sized SiC = 10/90 were
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M. Fukushima et al. / Materials Science and Engineering B 148 (2008) 211–214
blended in ethanol for 1 h by using a planetary mill with a SiC pot and SiC balls. After mixing, the slurry was dried with a rotary evaporator and a vacuum oven at 110 ◦ C for 4 h, and screened through a 110 mesh sieve. The powder compacts were formed to a disk shape without any binder by using a steel mold at 40 MPa, and they were then subjected to cold isostatic pressing (CIP) at either 100 or 400 MPa (hereafter identified as CIP100 or CIP400, respectively). The size of the formed disk specimens was approximately 20 mm in diameter and 5 mm in thickness. The green bodies were placed in a capped graphite crucible, and sintered at 1500, 1600, 1700 and 1800 ◦ C for 2 h under Ar gas flow. The heating rate from room temperature to 1200 ◦ C was 20 ◦ C/min, and 15 ◦ C/min from 1200 ◦ C to setting temperatures. The apparent densities of the sintered specimens were measured by water displacement using the Archimedes method, according to the equation [9]: d=
ρWa , Ww − W l
where Wl , Ww and Wa indicate the weight in water, wet weight and dry weight, respectively, and ρ is the water density at the measurement temperature. Green body densities were calculated from the dimensions and weight of the specimens. Specific surface area was measured by nitrogen adsorption/desorption (Yuasa Ionics Inc., Autosorb, Osaka, Japan), where surface area was determined from a BET (Brunauer, Emmet and Teller) analysis in the P/P0 range of 0.05–0.30 using a molecular cross-sectional area for N2 of 0.163 nm2 and 10 points [10]. The microstructures were observed using a scanning electron microscopy (SEM; JEOL-6330F, Japan). The pore size distributions were measured by mercury porosimetry (Yuasa Ionics Inc., PoreMaster-GT, Osaka, Japan) in the range of 6.1 nm–426 m. The weight of the specimens used for the measurement was around 0.5 g. The compressive strength was measured according to JIS R1608. The used cylindrical specimen had a dimension of 5 mm and 12.5 mm in height and was fractured using a testing machine (MTS Systems Corporation, Sintech 10/GL, Minnesota, USA). The load was applied on the two parallel surfaces at a cross-head speed of 0.5 mm/min. For every material, five specimens were tested.
Table 1 Linear shrinkage and relative density vs. sintering temperature CIP pressure
Sintering temperature (◦ C)
Linear shrinkage (%)
Relative density (%)
100 MPa
Green body 1500 1600 1700 1800
– 0.5 0.5 0.5 0.5
55 56 56 56 56
400 MPa
Green body 1500 1600 1700 1800
– 0.2 0.2 0.2 0.2
60 61 61 61 61
Fig. 1 shows typical examples of the pore size distributions of the specimens treated with different CIP pressures and sintered at different temperatures. The pore sizes of the specimens sintered at 1500 ◦ C were in the ranges of 0.10–0.22 and 0.09–0.19 m for CIP100 and CIP400 specimens, respectively. When sintered at 1800 ◦ C, the pore sizes increased to 0.50–1.15 m for CIP100 and 0.40–0.94 m for CIP400, respectively. Hence, higher CIP pressure reduced the pore size and high sintering temperature resulted in enlarged pore sizes. The peak pore size shown in Fig. 2 demonstrates a clear relationship between pore size, sintering temperature, nano-sized powder addition, and CIP pressure. The peak sizes for CIP100 were 0.19, 0.50, 0.75 and 0.95 m from 1500 to 1800 ◦ C, and those for CIP400 were 0.14, 0.41, 0.57 and 0.73 m, respectively. Table 2 shows the typical values of BET surface area and compressive strength of the specimens treated at CIP400 and sintered at 1500 and 1800 ◦ C. The specimens sintered at 1500 ◦ C had a surface area of 9.3 m2 /g. However, when sintered at 1800 ◦ C, the surface area decreased to 1.9 m2 /g, which was about 20% of that of the specimen sintered at 1500 ◦ C. Since the porosity measured by the Archimedes method was constant during sintering, the decrease in surface area is considered to be due to the decrease of the number of fine particles, grain growth
3. Results and discussion The linear shrinkage and relative density as a function of sintering temperature are listed in Table 1. The specimens did not exhibit significant linear shrinkage, regardless of sintering temperature. An effect of CIP pressure on density was clearly observed; higher pressure resulting in higher density. For all sintering temperatures, the specimens treated at CIP400 MPa exhibited 5% higher relative density than those treated at CIP100 MPa. In addition, densities were constant regardless of sintering temperature. These observed results are consistent with previous reports of -SiC compacts produced without any sintering aids [6,11]. Hase et al. have reported that heating a compact of -SiC with an average particle size of 0.1 m to 2100 ◦ C resulted in no densification or shrinkage.
Fig. 1. Pore size distributions of porous SiC sintered at 1500 and 1800 ◦ C.
M. Fukushima et al. / Materials Science and Engineering B 148 (2008) 211–214
Fig. 2. Peak pore size of porous SiC sintered at various temperatures. Table 2 Typical values of BET surface area and compressive strength of the specimens treated at CIP400 and sintered at 1500 and 1800 ◦ C Sintering temperature (◦ C)
BET surface area (m2 /g)
Compressive strength (MPa)
1500 1800
9.3 ± 0.6 1.9 ± 0.2
122 ± 22 513 ± 38
Fig. 3. SEM micrographs of the fractured surfaces of the CIP400 specimens sintered at (a) 1500 and (b) 1800 ◦ C, respectively.
213
and enlarged pore size. The results of strength tests showed that the compressive strength of the specimens sintered at 1500 ◦ C was 122 MPa, and this increased to 522 MPa when sintering at 1800 ◦ C. Although the pore size of the specimen sintered at 1800 ◦ C was 5 times larger than that sintered at 1500 ◦ C, the increase of sintering temperature resulted in increased strength, suggesting that well-developed neck areas were formed at higher sintering temperatures. Fig. 3 shows SEM micrographs of the fractured surfaces of the CIP400 specimens sintered at (a) 1500 and (b) 1800 ◦ C, respectively. It is clearly seen that the microstructure changed with sintering temperature. Fine particles and small pores were observed in the specimens sintered at 1500 ◦ C but disappeared after sintering at 1800 ◦ C, where appreciable grain growth and enlarged pores were observed. These phenomena can be more clearly seen in high magnification observations, as shown in Fig. 4. In addition, large necks were observed in the specimen sintered at 1800 ◦ C. In general, the surface of the SiC particles is covered with a thin, naturally existing, SiO2 layer. When the green compact is heated above 1400 ◦ C, this surface SiO2 layer can react with free carbon (one of the main impurities) or SiC, resulting in
Fig. 4. SEM micrographs of the fractured surfaces of the CIP400 specimens sintered at (a) 1500 and (b) 1800 ◦ C, respectively.
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the evolution of SiO and CO gas, according to the following equations [4,11–14]: SiO2 + C → SiO(g) + CO(g) 2SiO2 + SiC → 3SiO(g) + CO(g) SiO(g) + 2C(s) → SiC(s) + CO(g) Through these reactions a silica free surface is formed on the SiC particles, which can accelerate the surface diffusion. The diffusion process from the particle surface to the neck regions should enlarge the neck area. This may result in the consumption of the nano-sized particles, since they have larger surface energy and smaller diffusion distance than submicron-sized particles. Mass transfer from the surface towards the neck should result in an increase in both pore size and grain size as well as the neck area. Compared to our previous reports, the neck area observed in this study was well developed, which was due to the promoted mass transfer by nano-sized powder additions [15]. The present compressive strength (513 MPa) was also much higher than the strength reported previously (15 MPa) for our previous porous SiC with similar density (58%) and peak pore size (1.7 m). Thus, this study suggests that fabricating porous SiC by promoting mass transfer can provide high strength due to increased neck area. 4. Conclusion Porous silicon carbide sintered at 1500–1800 ◦ C was prepared by using different CIP pressures and nano-sized SiC additives. The role of these processing factors on pore size and microstructure was investigated. A higher CIP pressure treatment was effective in reducing pore size. The pore size and particle size increased with increasing sintering temperature, and
the disappearance of the added nano-sized powder was observed. Enlarged neck regions were also found, due to mass transfer of the nano-sized particles from the particle surface toward the neck area. The compressive strength increased with sintering temperature and attained a value over 500 MPa through the formation of well-developed neck regions. The addition of nano-sized powder suggests that the promoted mass transfer from the surface to neck regions can provide high strength due to increased neck area. References [1] M.J. Ledoux, S. Hantzer, C.P. Huu, J. Guille, M.P. Desaneaux, J. Catal. 114 (1988) 176–185. [2] M. Benaissa, J. Werckmann, G. Ehret, E. Peschiera, J. Guille, M.J. Ledoux, J. Mater. Sci. 29 (1994) 4700–4707. [3] N. Keller, C. Pham-Huu, S. Roy, M.J. Ledoux, C. Estournes, J. Guille, J. Mater. Sci. 34 (1999) 3189–3202. [4] V. Suwanmethanond, E.P. Goo, K.T. Liu, G. Johnston, M. Sahimi, T.T. Tsotsis, Ind. Eng. Chem. Res. 39 (2000) 3264–3271. [5] P.-K. Lin, D.-S. Tsai, J. Am. Ceram. Soc. 80 (1997) 365–372. [6] M. Fukushima, Y. Zhou, Y. Iwamoto, S. Yamazaki, T. Nagano, H. Miyazaki, Y. Yoshizawa, K. Hirao, J. Am. Ceram. Soc. 89 (2006) 1523–1529. [7] J. She, J.-F. Yang, N. Kondo, T. Ohji, S. Kanzaki, Z.-Y. Deng, J. Am. Ceram. Soc. 85 (2002) 2852–2854. [8] F. Wakai, M. Yoshida, Y. Shinoda, T. Akatsu, Acta Mater. 53 (2005) 1361–1371. [9] J.T. Jones, M.F. Berard, Ceramics—Industrial Processing and Testing, The Iowa State University Press, Ames, Iowa, 1993, pp. 172–174. [10] S. Brunauer, P.H. Emmett, E. Teller, J. Am. Chem. Soc. 60 (1938) 309–319. [11] T. Hase, H. Suzuki, I. Tomozuka, Yogyo-Kyoukai-shi 87 (1979) 317–321 (in Japanese). [12] C. Greskovich, J.H. Rosolowski, J. Am. Ceram. Soc. 59 (1976) 336–343. [13] T. Hase, H. Suzuki, Yogyo-Kyoukai-shi 88 (1980) 258–264 [in Japanese]. [14] T. Grande, H. Sommerset, E. Hagen, K. Wiik, M.-A. Einarsrud, J. Am. Ceram. Soc. 80 (1997) 1047–1052. [15] M. Fukushima, Y. Zhou, Y. Yoshizawa, H. Miyazaki, K. Hirao, Presented at the 31st International Cocoa Beach Conference & Exposition on Advanced Ceramics & Composites, 2007.