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Effects of B4C on the microstructure and phase transformation of porous SiC ceramics
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CERAMICS INTERNATIONAL
Ceramics International ] (]]]]) ]]]–]]] www.elsevier.com/locate/ceramint
Effects of B4C on the microstructure and phase transformation of porous SiC ceramics Wenming Guoa,b, Hanning Xiaoc,n, Jingxiong Liuc, Jianjun Liangc, Pengzhao Gaoc, Guangming Zenga,b b
a College of Environmental Science and Engineering, Hunan University, Changsha 410082, China Key Laboratory of Environmental Biology and Pollution Control (Hunan University), Ministry of Education, Changsha 410082, China c College of Materials Science and Engineering, Hunan University, Changsha 410082, China
Received 29 March 2015; received in revised form 11 May 2015; accepted 12 May 2015
Abstract Porous SiC ceramics sintered with B4C additive exhibit excellent corrosion resistance in hot acidic and basic solutions compared with traditional porous ceramic membrane supports. The effects of B4C on the density, phase composition, and microstructure of porous 6H–SiC ceramics were investigated in this study. The phase composition and microstructure of the as-prepared samples were characterized through X-ray diffraction, Raman spectrum, scanning electron microscopy, and electron probe microanalyzer equipped with wavelength dispersive spectrometry. B4C promoted the sintering of porous SiC ceramics within the temperature range of 2100–2200 1C. The B4C content higher than 0.5 wt% induced the phase transformation of 6H–SiC to 4H–SiC at 2200 1C, resulting in the formation of large plate-like 4H–SiC grains, which significantly improved the flexural strength of the porous ceramics. The mechanisms of phase transformation and grain formation were also elucidated. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved. Keywords: A. Sintering; B. Microstructure final; D. SiC; E. Membranes; Porous ceramics
1. Introduction Ceramic membranes are applied in different industries because of their remarkable properties, such as high separation efficiency, excellent thermal stability, high-pressure resistance, chemical stability, and long service life, compared with polymeric membranes [1–3]. Studies and application data have shown that ceramic membranes present some limitations when exposed in harsh environments, such as hot corrosive solutions; among which, gradual degradation of the mechanical property of porous membrane supports may result in the damage of membranes under operating pressures [4–7]. It suggests that the main phases in the support such as Al2O3, or the inter-granular phase used to facilitate the sintering of porous support, cannot tolerate acid or alkali corrosion. n
Corresponding author. Tel.: þ86 731 88822269; fax: þ 86 731 88823554. E-mail address: [email protected] (H. Xiao).
Covalently bonded materials, such as SiC, have been considered as substitutes for traditional support materials to improve the corrosion resistance of the support [8,9]. Recrystallized SiC sintered without any additives can endure acidic and alkaline solution and high temperatures and has been used to support commercial ceramic membranes under severe conditions [10]. However, the high cost of recrystallized SiC associated with the high sintering temperature requirement (4 2300 1C) and its relatively low strength have restricted its wide applications. Porous SiC supports can be fabricated through reaction bonding [11,12] and liquid-phase sintering [13,14] at low sintering temperatures. Investigations on the behavior of dense reaction-bonded SiC and liquid-phase-sintered SiC in corrosive solutions have revealed that residual Si and sintering additives are detrimental to the corrosion resistance of SiC ceramics [15–18]. Kim et al. reported that the residual free Si in reaction-bonded SiC is preferentially corroded in water at 360 1C [15]. Mikeska et al. stated that reaction-bonded SiC shows poor corrosion resistance in HF solution at 90 1C
http://dx.doi.org/10.1016/j.ceramint.2015.05.059 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059
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because of free Si [16]. Ray et al. evaluated the corrosion of SiC sintered with Al or AlN in HF at 80 1C and confirmed that the grain boundary phase is vulnerable to attack, which resulted in severe strength degradation [17]. Ray et al. further found that SiC ceramics sintered with 2.5 wt% AlN and subjected to different heating rates exhibit varied corrosion resistance [18]. Thus, these findings showed the significant influence of changes in the amount or composition of the grain boundary phase. We previously fabricated a SiC porous support through solidstate sintering by using B4C, and the resultant support maintained excellent corrosion resistance in hot acidic and alkaline solutions [19]. Interestingly, porous ceramics sintered with more than 1 wt% B4C at 2200 1C presented a microstructure with equiaxed SiC particles interpenetrated with large plate-like grains; these ceramics also demonstrated an approximately 50% increase in strength compared with those without plate-like grains. Nevertheless, the mechanism underlying the formation of plate-like grains, which can be the transformation of 4H–SiC from 6H–SiC, remains unclear. Large plate-like grains have been observed in SiC ceramics, which endure a phase transformation during sintering or annealing [20–29]. The formation of plate-like grains may improve the toughness of SiC ceramics by introducing crystal seeds in the starting materials, such as in β-SiC (3C–SiC) [20–22]. Heuer et al. investigated three annealed SiC ceramics containing 3C–SiC and showed that plate-like α-SiC was formed because of the preferential growth of the low-energy coherent β/α interfaces, thereby producing α(core)/β(envelope) composite grains at the initial stage of β-α transformation [23,24]. Malinge et al. performed pressureless sintering of β-SiC nanoparticles to support Heuer's theory; in this study, α-SiC growth was initiated from stacking faults initially present in β-SiC and required β-SiC recrystallization to create the {111}β || (0001)α interfaces, which determine the transformation rate [25]. Seo et al. also found that the 3C-4H phase transformation originated at the stacking faults and accelerated by adding Al [26]. Aluminum impurities in SiC raw materials can also stabilize 4H–SiC and accelerate 6H–SiC transformation, resulting in the formation of elongated grains [27–29].
Fig. 2. XRD spectra of porous SiC ceramics with different B4C contents sintered at different temperatures.
The appearance of plate-like grains and phase transformation in our study differed from that of previous findings [19]. The temperature for the plate-like grain formation and phase transformation was higher than that previously reported, and the starting materials did not contain any other polytype SiC as the seeds of plate-like grains. Moreover, the amount of Al impurity in the starting materials may be insufficient to accelerate phase transformation as evidenced by the undetected transformation of 6H–SiC to 4H–SiC in batches with low B4C contents,. Hence, the present work aims to explain this phenomenon and reveal the relationship between plate-like grains and phase transformation of 6H4H SiC. 2. Experimental 2.1. Raw materials Coarse 6H–SiC powder ( 28 μm) and submicron 6H–SiC powder ( 0.5 μm) were obtained from Changle Xinyuan Carborundum Micropowder Co., Ltd. (Shandong, China). B4C powder ( 1.0 μm, Jingangzuan Boron Carbide Co., Ltd., China) and polyvinyl alcohol (PVA, average polymerization degree of 1750750, Sino Chemical Reagent Co., Ltd., China) were used as sintering agent and temporary binder, respectively. 2.2. Preparation
Fig. 1. Density of porous SiC ceramics sintered at different temperatures. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
The weight ratio of coarse SiC, submicron SiC, and PVA was 70:30:1. The B4C contents of SiC0, SiC0.25, SiC0.5, SiC1, SiC2, and SiC3 were 0, 0.25, 0.5, 1, 2, and 3% by weight ratio, respectively. The starting powder mixtures (SiC and B4C) were blended through ball milling in 1 wt% PVA solution with SiC grinding balls in a polyurethane jar (400 r/min, 1 h) and then dried by rotary evaporation. Subsequently, the powder mixtures were sieved with a 60-mesh screen and pressed under 30 MPa. The compacts were loaded into recrystallized SiC crucibles and individually sintered at 1900, 2100, 2150, and 2200 1C for 1 h.
Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059
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Fig. 3. Raman spectra of porous SiC ceramics with different B4C contents sintered at 2200 oC. Table 1 Raman frequencies of fundamental SiC polytypes30. Polytype
4H 6H
Frequency (cm 1) Planar acoustic FTA
Planar optic FTO
196, 204, 266 145, 150, 236, 241, 266
796, 776 797, 789
identification was performed through powder X-ray diffraction (XRD, Rigaku D/max2200 VPC) by using monochromic Cu Kα radiation of 0.15405 nm operated at 40 kV and 40 mA and Raman spectroscopy (Renishaw invia-reflex) by using a laser with 633 nm wavelength. The morphology and microstructures of the samples were observed under a scanning electron microscope (JEOL, JSM6490LV). Elemental analysis of porous ceramics was performed using an electron probe microanalyzer equipped with a wavelength-dispersive X-ray spectroscopy (JEOL, JXA-8230).
2.3. Characterization
3. Results and discussion
The density of the samples was measured through the water displacement technique (Archimede's method). Flexural strength was determined using the three-point bending method in a span of 35 mm at room temperature and loading rate of 0.5 mm/min for samples with a geometrical size of 8 mm 6 mm 65 mm. Phase
3.1. Effects of B4C on the density of porous SiC ceramics The variation in the density of porous SiC ceramics at different sintering temperatures is shown in Fig. 1. The density of the samples treated at 1900 1C for 1 h is designated as the density of the green
Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059
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10μm
10μm
Fig. 4. Fracture surface of SiC0 sintered at 2100 1C (a) and 2200 1C (b).
10μm Fig. 5. Fracture surface of SiC1 sintered at 2100 1C.
body because the samples showed no linear shrinkage during heat treatment. Since the samples can be subjected to water displacement technology without incurring any damage, they can be directly compared with those sintered at higher temperatures. In this method, the density difference caused by the methods used to measure three dimensions and water displacement can be avoided. As shown in Fig. 1, the variation in the green density showed no regularity with B4C content and the maximum difference of the green density for different formula was 0.03 g/cm3. This finding can be attributed to the different residual water contents in the sieved powder, which affected the subsequent pressing process. However, the standard deviation was lower than 0.009 g/cm3 for each batch. We can study the effect of B4C on the density of porous SiC ceramics by using the remainder between the sintered and green densities. For SiC0 without B4C additive, the density stabilized after sintering from 2100 1C to 2200 1C. This character is similar to that of recrystallized SiC, which resulted from the non-densifying sintering mechanisms of surface diffusion and evaporation–condensation. For those doped with B4C, the density of porous SiC ceramics increased with temperature. The increase was weak (o 0.02 g/cm3) when the B4C content was lower than 1% but more than 0.04 g/cm3 and 0.06 g/ cm3 at 2100 1C and 2200 1C, respectively, for B4CZ1 wt%. The increase in density for different batches was insignificant when the B4C content was higher than 1 wt%. 3.2. Effects of B4C on the phase composition of porous SiC ceramics The XRD patterns of porous SiC ceramics sintered at different temperatures with different B4C contents are shown in Fig. 2. Only 6H–SiC was detected in samples sintered at
temperatures lower than 2200 1C or those with o 1% B4C at 2200 1C for 1 h. The peaks of 4H–SiC could be observed when sintered at 2200 1C for SiC1–SiC3. The peak intensity of 4H–SiC increased with increasing B4C content, whereas that of 6H–SiC decreased. Since the Raman scattering efficiency of SiC is relatively strong and signals with a high signal/noise ratio can be collected in a short time, Raman scattering spectroscopy has been used to identify polytype structures of SiC [30–32]. In addition, the folded modes of the transverse optic (FTO) and acoustic phonon (FTA) branches for different SiC polytypes are separately located. Therefore, the FTO and FTA modes are used to characterize SiC polytypes when mixtures of different polytypes are present. The Raman spectra of porous SiC0.5–3 sintered at 2200 1C for 1 h in the FTA and FTO regions are shown in Fig. 3, and the Raman frequencies of the fundamental SiC polytypes are displayed in Table 1. For SiC0.5, all Raman bands corresponded to 6H–SiC. Increasing B4C content resulted in weaker Raman intensity for 6H–SiC in the FTA region (Fig. 3a), whereas stronger for 4H–SiC. This finding showed that 4H–SiC phase could be detected for B4C content Z 1%, and 4H–SiC content in porous ceramics increased with increasing B4C content. The Raman bands in the FTO region broadened when the B4C contentZ 1% as indicated by the bands for 4H–SiC that were close to those of 6H–SiC. However, the evolution of the shoulder peak near 775 cm 1 indicated that 4H–SiC content increased in SiC1–3 with increasing B4C content (Fig. 3b). Hence, the phase evolution observed in the Raman spectrum coincided with the XRD results and B4C considerably affected the phase transformation of 6H–SiC to 4H–SiC. 3.3. Effects of B4C on the microstructure of porous SiC ceramics Fig. 4 shows the microstructure of SiC0 sintered at 2100 1C and 2200 1C. The starting submicron SiC particles disappeared, and several fractured necks and large particles ( 30 μm) with round morphology and connected with weak necks were observed. For SiC1 sintered at 2100 1C (Fig. 5), particles maintained their morphology similar to SiC0 but particles in SiCl were connected with stronger necks. As fractures were formed at necks and in the particles, the higher strength of SiC1 sintered at 2100 1C than that of SiC0 (Fig. 4a) can be attributed to stronger necks. When the
Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059
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50μm
10μm
Fig. 6. Fracture surface of SiC0.5 sintered at 2200 1C: (a) area with equiaxed grains (b) area with plate-like grains.
100μm
100μm Fig. 7. Fracture surface of SiC1 (a) and SiC 3 (b) sintered at 2200 1C.
Fig. 8. Flexural strength of porous SiC ceramics with or without B4C additives sintered at 2100–2200 1C.
temperature reached 2200 1C, the SiC0.5 microstructure predominantly exhibited equiaxed particles (Fig. 6a) but several large platelike grains ( 100 μm in length) could be occasionally observed (Fig. 6b). With increasing B4C content (Fig. 7a and b), the quantity of the plate-like grains increased gradually but their sizes increased initially and then decreased. The plate-like grains can interpenetrate with several equiaxed particles (either combined or uncombined) to form a strong bonded network. The bonding neck of the plate-like grain and nearby particles was almost as large as the particle. The microstructure of porous SiC ceramics with interpenetrated
network and strong bonding facilitated strength improvement (Fig. 8). Moreover, the strength of SiC1 was approximately 50% higher than those with the equiaxed microstructure. The microstructure also affected the spreading of the crack. Fig. 9a shows that the samples with the equiaxed microstructure exhibited a zigzag surface, whereas those with interpenetrated network and strong bonding showed a planar fracture surface (Fig. 9b). The forementioned observations on the microstructure evolution of porous SiC ceramics indicated that the formation of plate-like grains occurred at 2200 1C and Z0.5% B4C content. As shown in Fig. 10, four local analyses of Raman scattering technique demonstrated that the plate-like grains were 4H–SiC. In addition to the phase evolution revealed by XRD and Raman analyses, the abnormal growth of plate-like 4H–SiC grains seemed simultaneous with the phase transformation of 6H to 4H polytype. It should be noted that the 4H–SiC was not detected by XRD and Raman analyses for SiC0.5 sintered at 2200 1C was because the content exceeded the detection threshold of the two techniques. 3.4. Mechanism of phase transformation of 6H–SiC and grain growth of plate-like 4H–SiC Porous SiC samples containing B4C shrank at 2100 1C to 2200 1C, indicating that B4C promoted SiC sintering. Different B4C contents resulted in varied shrinkage degrees, i.e., shrinkage increased when the B4C content was o 1% and became stable when B4C 4 1%. Plate-like 4 H–SiC grains emerged when B4C reached 0.5%, and its amount increased with increasing B4C content. The evolution of sintering and
Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059
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500μm
500μm
Fig. 9. Crack propagation of porous ceramics: (a) SiC0.25 and (b) SiC3 sintered at 2200 1C for 1 h.
Fig. 10. FTA bands of Raman shifts for the plate-like grains in SiC3 sintered at 2200 1C.
microstructure can be interpreted by the solid solution of B4C in SiC. Numerous researchers have investigated in detail the effects of B/B4C and C additives on the solid-state sintering of dense SiC ceramics and considered that the B/B4C solutes in SiC created vacancies of silicon and carbon, which enhanced lattice diffusion and caused compaction shrinkage [33,34]. Similarly, the density variation of porous SiC ceramics as B4C can be explained. The solubility of B in SiC is approximately 0.3 wt% at 2040 1C [33], and a slightly higher solubility limit of B4C is 0.5 wt% at 2200 1C [35]. In the present work, the original large SiC particles with low sintering activity determined the minimal increase in density, which is beneficial in the preparation of porous ceramics. However, B4C doping in SiC in solid solution enhanced the diffusion rate of Si and C, which led to the mass transfer of surface diffusion in pure SiC
substituted by the lattice diffusion for those containing B4C as a result of density increase. For SiC containing B4C content lower than the solubility limit, high B4C content may generate high impurity concentration, resulting in high diffusion rate. Consequently, the density of SiC porous ceramics increased with increasing B4C content and then maintained a steady state. The formation of large plate-like 4H–SiC grains is related to the content and distribution sites of B4C in porous SiC ceramics. Based on the microstructure observation and the phase evolution of SiC, the nucleation and growth of plate-like 4H–SiC grains were dependent on excessive B4C in local regions. 4H–SiC nucleation may originate at the SiC particles combined with B4C particles, resulting in excessive B4C in the local regions. The local regularity of 6H–SiC became disordered because of
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improved the flexural strength. The phase transformation and formation of large plate-like grains were associated with locally excessesive B4C, which exceeded the solubility limit of B4C in 6H–SiC, resulting in the formation of stacking faults that provide 4H–SiC nucleus. Moreover, the diffusion of C and Si atoms from 6H–SiC to 4H–SiC nucleus through the difference of polytype stability and B4C concentration resulted in the plate-like grain growth of 4H–SiC and phase transformation. Acknowledgments The authors are grateful for the financial support from the National Natural Science Foundation of China (Grant nos. 51302076, 51372078) and the China Postdoctoral Science Foundation (Grant no. 2013M531783). Fig. 11. Element composition of the plate-like grain in SiC3 sintered at 2200 1C for 1 h.
high B4C concentration, which induced the unstable state of 6H–SiC and formed defects. The broad FTO bands of Raman shift for 6H–SiC doped with B4C in Fig. 3b indicated heavy stacking disorders and defects, which were also observed by Nakashima et al. [30]. Based on this finding, some stacking fault may be a suitable displacement layer of 6H–SiC to form the initial stage of an original 4H–SiC nucleus. 4H–SiC with fewer periodic stacking sequences of Si–C close-packed atomic planes seemed more stable than 6H–SiC when doped with other elements; this observation is similar to the finding of Takana et al. [27,28]. Si and C atoms diffused from 6H–SiC to the nucleus containing abundant B to minimize free energy and stabilize the structure, and this diffusion led to 4H–SiC grain growth. As the growth of the newly formed 4H–SiC grain contained fewer defects, i.e., with lower B concentration, B atoms diffused in the opposite direction of Si and C atoms. Fig. 11 shows that this change resulted in higher B concentration in the region around 4H–SiC than that in the center. Hence, the absence of plate-like 4H–SiC grains in the SiC0.25 microstructure was due to the lack of locally excessive B4C in which the B4C content was lower than the limit of solid solubility. For SiC0.5 with B4C content near the solubility limit, sites with excessive B4C content may locally exist. This excess resulted in occasional formation of plate-like 4H–SiC grains. Moreover, the quantity of plate-like 4H–SiC grains increased with increasing B4C content for SiC1–3 because of increasing excess B4C sites. 4. Summary B4C slightly promoted sintering of porous SiC ceramics between 2100 1C and 2200 1C, which maintained the porosity of the samples. The transformation of 6H–SiC to 4H–SiC was induced by B4C at 2200 1C, which was accompanied by the appearance of large plate-like 4H–SiC grains interpenetrating with several equiaxed particles; these processes significantly
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CERAMICS INTERNATIONAL
Ceramics International ] (]]]]) ]]]–]]] www.elsevier.com/locate/ceramint
Effects of B4C on the microstructure and phase transformation of porous SiC ceramics Wenming Guoa,b, Hanning Xiaoc,n, Jingxiong Liuc, Jianjun Liangc, Pengzhao Gaoc, Guangming Zenga,b b
a College of Environmental Science and Engineering, Hunan University, Changsha 410082, China Key Laboratory of Environmental Biology and Pollution Control (Hunan University), Ministry of Education, Changsha 410082, China c College of Materials Science and Engineering, Hunan University, Changsha 410082, China
Received 29 March 2015; received in revised form 11 May 2015; accepted 12 May 2015
Abstract Porous SiC ceramics sintered with B4C additive exhibit excellent corrosion resistance in hot acidic and basic solutions compared with traditional porous ceramic membrane supports. The effects of B4C on the density, phase composition, and microstructure of porous 6H–SiC ceramics were investigated in this study. The phase composition and microstructure of the as-prepared samples were characterized through X-ray diffraction, Raman spectrum, scanning electron microscopy, and electron probe microanalyzer equipped with wavelength dispersive spectrometry. B4C promoted the sintering of porous SiC ceramics within the temperature range of 2100–2200 1C. The B4C content higher than 0.5 wt% induced the phase transformation of 6H–SiC to 4H–SiC at 2200 1C, resulting in the formation of large plate-like 4H–SiC grains, which significantly improved the flexural strength of the porous ceramics. The mechanisms of phase transformation and grain formation were also elucidated. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved. Keywords: A. Sintering; B. Microstructure final; D. SiC; E. Membranes; Porous ceramics
1. Introduction Ceramic membranes are applied in different industries because of their remarkable properties, such as high separation efficiency, excellent thermal stability, high-pressure resistance, chemical stability, and long service life, compared with polymeric membranes [1–3]. Studies and application data have shown that ceramic membranes present some limitations when exposed in harsh environments, such as hot corrosive solutions; among which, gradual degradation of the mechanical property of porous membrane supports may result in the damage of membranes under operating pressures [4–7]. It suggests that the main phases in the support such as Al2O3, or the inter-granular phase used to facilitate the sintering of porous support, cannot tolerate acid or alkali corrosion. n
Corresponding author. Tel.: þ86 731 88822269; fax: þ 86 731 88823554. E-mail address: [email protected] (H. Xiao).
Covalently bonded materials, such as SiC, have been considered as substitutes for traditional support materials to improve the corrosion resistance of the support [8,9]. Recrystallized SiC sintered without any additives can endure acidic and alkaline solution and high temperatures and has been used to support commercial ceramic membranes under severe conditions [10]. However, the high cost of recrystallized SiC associated with the high sintering temperature requirement (4 2300 1C) and its relatively low strength have restricted its wide applications. Porous SiC supports can be fabricated through reaction bonding [11,12] and liquid-phase sintering [13,14] at low sintering temperatures. Investigations on the behavior of dense reaction-bonded SiC and liquid-phase-sintered SiC in corrosive solutions have revealed that residual Si and sintering additives are detrimental to the corrosion resistance of SiC ceramics [15–18]. Kim et al. reported that the residual free Si in reaction-bonded SiC is preferentially corroded in water at 360 1C [15]. Mikeska et al. stated that reaction-bonded SiC shows poor corrosion resistance in HF solution at 90 1C
http://dx.doi.org/10.1016/j.ceramint.2015.05.059 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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because of free Si [16]. Ray et al. evaluated the corrosion of SiC sintered with Al or AlN in HF at 80 1C and confirmed that the grain boundary phase is vulnerable to attack, which resulted in severe strength degradation [17]. Ray et al. further found that SiC ceramics sintered with 2.5 wt% AlN and subjected to different heating rates exhibit varied corrosion resistance [18]. Thus, these findings showed the significant influence of changes in the amount or composition of the grain boundary phase. We previously fabricated a SiC porous support through solidstate sintering by using B4C, and the resultant support maintained excellent corrosion resistance in hot acidic and alkaline solutions [19]. Interestingly, porous ceramics sintered with more than 1 wt% B4C at 2200 1C presented a microstructure with equiaxed SiC particles interpenetrated with large plate-like grains; these ceramics also demonstrated an approximately 50% increase in strength compared with those without plate-like grains. Nevertheless, the mechanism underlying the formation of plate-like grains, which can be the transformation of 4H–SiC from 6H–SiC, remains unclear. Large plate-like grains have been observed in SiC ceramics, which endure a phase transformation during sintering or annealing [20–29]. The formation of plate-like grains may improve the toughness of SiC ceramics by introducing crystal seeds in the starting materials, such as in β-SiC (3C–SiC) [20–22]. Heuer et al. investigated three annealed SiC ceramics containing 3C–SiC and showed that plate-like α-SiC was formed because of the preferential growth of the low-energy coherent β/α interfaces, thereby producing α(core)/β(envelope) composite grains at the initial stage of β-α transformation [23,24]. Malinge et al. performed pressureless sintering of β-SiC nanoparticles to support Heuer's theory; in this study, α-SiC growth was initiated from stacking faults initially present in β-SiC and required β-SiC recrystallization to create the {111}β || (0001)α interfaces, which determine the transformation rate [25]. Seo et al. also found that the 3C-4H phase transformation originated at the stacking faults and accelerated by adding Al [26]. Aluminum impurities in SiC raw materials can also stabilize 4H–SiC and accelerate 6H–SiC transformation, resulting in the formation of elongated grains [27–29].
Fig. 2. XRD spectra of porous SiC ceramics with different B4C contents sintered at different temperatures.
The appearance of plate-like grains and phase transformation in our study differed from that of previous findings [19]. The temperature for the plate-like grain formation and phase transformation was higher than that previously reported, and the starting materials did not contain any other polytype SiC as the seeds of plate-like grains. Moreover, the amount of Al impurity in the starting materials may be insufficient to accelerate phase transformation as evidenced by the undetected transformation of 6H–SiC to 4H–SiC in batches with low B4C contents,. Hence, the present work aims to explain this phenomenon and reveal the relationship between plate-like grains and phase transformation of 6H4H SiC. 2. Experimental 2.1. Raw materials Coarse 6H–SiC powder ( 28 μm) and submicron 6H–SiC powder ( 0.5 μm) were obtained from Changle Xinyuan Carborundum Micropowder Co., Ltd. (Shandong, China). B4C powder ( 1.0 μm, Jingangzuan Boron Carbide Co., Ltd., China) and polyvinyl alcohol (PVA, average polymerization degree of 1750750, Sino Chemical Reagent Co., Ltd., China) were used as sintering agent and temporary binder, respectively. 2.2. Preparation
Fig. 1. Density of porous SiC ceramics sintered at different temperatures. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
The weight ratio of coarse SiC, submicron SiC, and PVA was 70:30:1. The B4C contents of SiC0, SiC0.25, SiC0.5, SiC1, SiC2, and SiC3 were 0, 0.25, 0.5, 1, 2, and 3% by weight ratio, respectively. The starting powder mixtures (SiC and B4C) were blended through ball milling in 1 wt% PVA solution with SiC grinding balls in a polyurethane jar (400 r/min, 1 h) and then dried by rotary evaporation. Subsequently, the powder mixtures were sieved with a 60-mesh screen and pressed under 30 MPa. The compacts were loaded into recrystallized SiC crucibles and individually sintered at 1900, 2100, 2150, and 2200 1C for 1 h.
Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059
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Fig. 3. Raman spectra of porous SiC ceramics with different B4C contents sintered at 2200 oC. Table 1 Raman frequencies of fundamental SiC polytypes30. Polytype
4H 6H
Frequency (cm 1) Planar acoustic FTA
Planar optic FTO
196, 204, 266 145, 150, 236, 241, 266
796, 776 797, 789
identification was performed through powder X-ray diffraction (XRD, Rigaku D/max2200 VPC) by using monochromic Cu Kα radiation of 0.15405 nm operated at 40 kV and 40 mA and Raman spectroscopy (Renishaw invia-reflex) by using a laser with 633 nm wavelength. The morphology and microstructures of the samples were observed under a scanning electron microscope (JEOL, JSM6490LV). Elemental analysis of porous ceramics was performed using an electron probe microanalyzer equipped with a wavelength-dispersive X-ray spectroscopy (JEOL, JXA-8230).
2.3. Characterization
3. Results and discussion
The density of the samples was measured through the water displacement technique (Archimede's method). Flexural strength was determined using the three-point bending method in a span of 35 mm at room temperature and loading rate of 0.5 mm/min for samples with a geometrical size of 8 mm 6 mm 65 mm. Phase
3.1. Effects of B4C on the density of porous SiC ceramics The variation in the density of porous SiC ceramics at different sintering temperatures is shown in Fig. 1. The density of the samples treated at 1900 1C for 1 h is designated as the density of the green
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10μm
10μm
Fig. 4. Fracture surface of SiC0 sintered at 2100 1C (a) and 2200 1C (b).
10μm Fig. 5. Fracture surface of SiC1 sintered at 2100 1C.
body because the samples showed no linear shrinkage during heat treatment. Since the samples can be subjected to water displacement technology without incurring any damage, they can be directly compared with those sintered at higher temperatures. In this method, the density difference caused by the methods used to measure three dimensions and water displacement can be avoided. As shown in Fig. 1, the variation in the green density showed no regularity with B4C content and the maximum difference of the green density for different formula was 0.03 g/cm3. This finding can be attributed to the different residual water contents in the sieved powder, which affected the subsequent pressing process. However, the standard deviation was lower than 0.009 g/cm3 for each batch. We can study the effect of B4C on the density of porous SiC ceramics by using the remainder between the sintered and green densities. For SiC0 without B4C additive, the density stabilized after sintering from 2100 1C to 2200 1C. This character is similar to that of recrystallized SiC, which resulted from the non-densifying sintering mechanisms of surface diffusion and evaporation–condensation. For those doped with B4C, the density of porous SiC ceramics increased with temperature. The increase was weak (o 0.02 g/cm3) when the B4C content was lower than 1% but more than 0.04 g/cm3 and 0.06 g/ cm3 at 2100 1C and 2200 1C, respectively, for B4CZ1 wt%. The increase in density for different batches was insignificant when the B4C content was higher than 1 wt%. 3.2. Effects of B4C on the phase composition of porous SiC ceramics The XRD patterns of porous SiC ceramics sintered at different temperatures with different B4C contents are shown in Fig. 2. Only 6H–SiC was detected in samples sintered at
temperatures lower than 2200 1C or those with o 1% B4C at 2200 1C for 1 h. The peaks of 4H–SiC could be observed when sintered at 2200 1C for SiC1–SiC3. The peak intensity of 4H–SiC increased with increasing B4C content, whereas that of 6H–SiC decreased. Since the Raman scattering efficiency of SiC is relatively strong and signals with a high signal/noise ratio can be collected in a short time, Raman scattering spectroscopy has been used to identify polytype structures of SiC [30–32]. In addition, the folded modes of the transverse optic (FTO) and acoustic phonon (FTA) branches for different SiC polytypes are separately located. Therefore, the FTO and FTA modes are used to characterize SiC polytypes when mixtures of different polytypes are present. The Raman spectra of porous SiC0.5–3 sintered at 2200 1C for 1 h in the FTA and FTO regions are shown in Fig. 3, and the Raman frequencies of the fundamental SiC polytypes are displayed in Table 1. For SiC0.5, all Raman bands corresponded to 6H–SiC. Increasing B4C content resulted in weaker Raman intensity for 6H–SiC in the FTA region (Fig. 3a), whereas stronger for 4H–SiC. This finding showed that 4H–SiC phase could be detected for B4C content Z 1%, and 4H–SiC content in porous ceramics increased with increasing B4C content. The Raman bands in the FTO region broadened when the B4C contentZ 1% as indicated by the bands for 4H–SiC that were close to those of 6H–SiC. However, the evolution of the shoulder peak near 775 cm 1 indicated that 4H–SiC content increased in SiC1–3 with increasing B4C content (Fig. 3b). Hence, the phase evolution observed in the Raman spectrum coincided with the XRD results and B4C considerably affected the phase transformation of 6H–SiC to 4H–SiC. 3.3. Effects of B4C on the microstructure of porous SiC ceramics Fig. 4 shows the microstructure of SiC0 sintered at 2100 1C and 2200 1C. The starting submicron SiC particles disappeared, and several fractured necks and large particles ( 30 μm) with round morphology and connected with weak necks were observed. For SiC1 sintered at 2100 1C (Fig. 5), particles maintained their morphology similar to SiC0 but particles in SiCl were connected with stronger necks. As fractures were formed at necks and in the particles, the higher strength of SiC1 sintered at 2100 1C than that of SiC0 (Fig. 4a) can be attributed to stronger necks. When the
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50μm
10μm
Fig. 6. Fracture surface of SiC0.5 sintered at 2200 1C: (a) area with equiaxed grains (b) area with plate-like grains.
100μm
100μm Fig. 7. Fracture surface of SiC1 (a) and SiC 3 (b) sintered at 2200 1C.
Fig. 8. Flexural strength of porous SiC ceramics with or without B4C additives sintered at 2100–2200 1C.
temperature reached 2200 1C, the SiC0.5 microstructure predominantly exhibited equiaxed particles (Fig. 6a) but several large platelike grains ( 100 μm in length) could be occasionally observed (Fig. 6b). With increasing B4C content (Fig. 7a and b), the quantity of the plate-like grains increased gradually but their sizes increased initially and then decreased. The plate-like grains can interpenetrate with several equiaxed particles (either combined or uncombined) to form a strong bonded network. The bonding neck of the plate-like grain and nearby particles was almost as large as the particle. The microstructure of porous SiC ceramics with interpenetrated
network and strong bonding facilitated strength improvement (Fig. 8). Moreover, the strength of SiC1 was approximately 50% higher than those with the equiaxed microstructure. The microstructure also affected the spreading of the crack. Fig. 9a shows that the samples with the equiaxed microstructure exhibited a zigzag surface, whereas those with interpenetrated network and strong bonding showed a planar fracture surface (Fig. 9b). The forementioned observations on the microstructure evolution of porous SiC ceramics indicated that the formation of plate-like grains occurred at 2200 1C and Z0.5% B4C content. As shown in Fig. 10, four local analyses of Raman scattering technique demonstrated that the plate-like grains were 4H–SiC. In addition to the phase evolution revealed by XRD and Raman analyses, the abnormal growth of plate-like 4H–SiC grains seemed simultaneous with the phase transformation of 6H to 4H polytype. It should be noted that the 4H–SiC was not detected by XRD and Raman analyses for SiC0.5 sintered at 2200 1C was because the content exceeded the detection threshold of the two techniques. 3.4. Mechanism of phase transformation of 6H–SiC and grain growth of plate-like 4H–SiC Porous SiC samples containing B4C shrank at 2100 1C to 2200 1C, indicating that B4C promoted SiC sintering. Different B4C contents resulted in varied shrinkage degrees, i.e., shrinkage increased when the B4C content was o 1% and became stable when B4C 4 1%. Plate-like 4 H–SiC grains emerged when B4C reached 0.5%, and its amount increased with increasing B4C content. The evolution of sintering and
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500μm
500μm
Fig. 9. Crack propagation of porous ceramics: (a) SiC0.25 and (b) SiC3 sintered at 2200 1C for 1 h.
Fig. 10. FTA bands of Raman shifts for the plate-like grains in SiC3 sintered at 2200 1C.
microstructure can be interpreted by the solid solution of B4C in SiC. Numerous researchers have investigated in detail the effects of B/B4C and C additives on the solid-state sintering of dense SiC ceramics and considered that the B/B4C solutes in SiC created vacancies of silicon and carbon, which enhanced lattice diffusion and caused compaction shrinkage [33,34]. Similarly, the density variation of porous SiC ceramics as B4C can be explained. The solubility of B in SiC is approximately 0.3 wt% at 2040 1C [33], and a slightly higher solubility limit of B4C is 0.5 wt% at 2200 1C [35]. In the present work, the original large SiC particles with low sintering activity determined the minimal increase in density, which is beneficial in the preparation of porous ceramics. However, B4C doping in SiC in solid solution enhanced the diffusion rate of Si and C, which led to the mass transfer of surface diffusion in pure SiC
substituted by the lattice diffusion for those containing B4C as a result of density increase. For SiC containing B4C content lower than the solubility limit, high B4C content may generate high impurity concentration, resulting in high diffusion rate. Consequently, the density of SiC porous ceramics increased with increasing B4C content and then maintained a steady state. The formation of large plate-like 4H–SiC grains is related to the content and distribution sites of B4C in porous SiC ceramics. Based on the microstructure observation and the phase evolution of SiC, the nucleation and growth of plate-like 4H–SiC grains were dependent on excessive B4C in local regions. 4H–SiC nucleation may originate at the SiC particles combined with B4C particles, resulting in excessive B4C in the local regions. The local regularity of 6H–SiC became disordered because of
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improved the flexural strength. The phase transformation and formation of large plate-like grains were associated with locally excessesive B4C, which exceeded the solubility limit of B4C in 6H–SiC, resulting in the formation of stacking faults that provide 4H–SiC nucleus. Moreover, the diffusion of C and Si atoms from 6H–SiC to 4H–SiC nucleus through the difference of polytype stability and B4C concentration resulted in the plate-like grain growth of 4H–SiC and phase transformation. Acknowledgments The authors are grateful for the financial support from the National Natural Science Foundation of China (Grant nos. 51302076, 51372078) and the China Postdoctoral Science Foundation (Grant no. 2013M531783). Fig. 11. Element composition of the plate-like grain in SiC3 sintered at 2200 1C for 1 h.
high B4C concentration, which induced the unstable state of 6H–SiC and formed defects. The broad FTO bands of Raman shift for 6H–SiC doped with B4C in Fig. 3b indicated heavy stacking disorders and defects, which were also observed by Nakashima et al. [30]. Based on this finding, some stacking fault may be a suitable displacement layer of 6H–SiC to form the initial stage of an original 4H–SiC nucleus. 4H–SiC with fewer periodic stacking sequences of Si–C close-packed atomic planes seemed more stable than 6H–SiC when doped with other elements; this observation is similar to the finding of Takana et al. [27,28]. Si and C atoms diffused from 6H–SiC to the nucleus containing abundant B to minimize free energy and stabilize the structure, and this diffusion led to 4H–SiC grain growth. As the growth of the newly formed 4H–SiC grain contained fewer defects, i.e., with lower B concentration, B atoms diffused in the opposite direction of Si and C atoms. Fig. 11 shows that this change resulted in higher B concentration in the region around 4H–SiC than that in the center. Hence, the absence of plate-like 4H–SiC grains in the SiC0.25 microstructure was due to the lack of locally excessive B4C in which the B4C content was lower than the limit of solid solubility. For SiC0.5 with B4C content near the solubility limit, sites with excessive B4C content may locally exist. This excess resulted in occasional formation of plate-like 4H–SiC grains. Moreover, the quantity of plate-like 4H–SiC grains increased with increasing B4C content for SiC1–3 because of increasing excess B4C sites. 4. Summary B4C slightly promoted sintering of porous SiC ceramics between 2100 1C and 2200 1C, which maintained the porosity of the samples. The transformation of 6H–SiC to 4H–SiC was induced by B4C at 2200 1C, which was accompanied by the appearance of large plate-like 4H–SiC grains interpenetrating with several equiaxed particles; these processes significantly
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Please cite this article as: W. Guo, et al., Effects of B4C on the microstructure and phase transformation of porous SiC ceramics, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.05.059