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Preparation of dense and high-purity SiC ceramics by pressureless solid-state-sintering
Ceramics International 45 (2019) 19771–19776
Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Preparation of dense and high-purity SiC ceramics by pressureless solidstate-sintering
T
Meng Liua,b, Yong Yanga,c,∗, Yuquan Weia, Yajie Lia,b, Heng Zhanga, Xuejian Liua,c, Zhengren Huanga,b,∗∗ a
Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China School of Physical Science and Technology, ShanghaiTech University, Shanghai, 201210, China c Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing, 100049, China b
ARTICLE INFO
ABSTRACT
Keywords: SiC Dense High-purity Pressureless solid-state-sintering
High-purity SiC materials have been used in semiconductor processes due to their excellent properties. However, they are difficult to densify without sintering aids. In this work, dense and high-purity SiC ceramics have been obtained by pressureless solid-state-sintering with ultra-low contents of sintering additives. The amount of residual B, C and O in the high-purity SiC ceramics was less than 0.15 wt%, respectively, and the total content of other impurity elements (such as aluminum, magnesium, calcium, iron, etc.) was less than 0.015 wt%. Finally, the purity of the as-prepared SiC ceramics was more than 99.5 wt%.
1. Introduction Silicon carbide (SiC) ceramics, due to strong covalent Si–C bond, possess many excellent properties such as high strength, good hardness, excellent chemical stability, outstanding oxidation and wear resistance [1–5]. Therefore, high-purity SiC ceramic devices are widely used in the various processing furnaces for silicon wafers (single crystal growth, oxidation, corrosion, ion injection, low-pressure CVD, etc.) [6]. Presently, silicon wafers with the diameter of 300 nm have been massively applied in IC manufacturing, and the wafers of 450 mm will be introduced in the near future due to the effect of Moore's Law. The larger wafers must be thicker to increase their resistance to warping and other structural deformities [7,8]. Therefore, there is a higher requirement of high-purity SiC ceramic devices with large size and high mechanical property used in the silicon wafer processing. Usually, large sized and complex shaped high-purity SiC ceramic devices can be fabricated by recrystallization sintering. Nevertheless, recrystallized silicon carbide (R–SiC) has lower flexural strength and poor oxidation resistance because of its relatively low density and porous structure. For example, Wenming Guo et al. [6] fabricated a high performance R–SiC with a density of 2.99 g/cm3 and flexural strength of 162.3 ± 2.6 MPa which was obtained by three polymer impregnation and pyrolysis-recrystallization cycles. But even so, the flexural strength of R–SiC can still hardly meet the requirement of the semiconductor industry.
∗
Many investigations have been done to increase the density of SiC ceramics without the sintering additive. Nadeau et al. [9] and Lara et al. [10] fabricated fully dense SiC bodies without sintering aids by hotpressed sintering at 2500 °C, 5000 MPa and spark plasma sintering at 2100 °C, 70 MPa, respectively. Nevertheless, neither of these methods is suitable for preparing SiC ceramic devices of large size and complex shapes. Pressureless solid-state-sintered SiC (S–SiC) could be dense and high-purity, and can be easily processed into ceramic devices with large size and complex shapes. Especially, it possesses a unique combination of properties including high thermal conductivity and outstanding hightemperature strength [11]. For S–SiC ceramics, boron-carbon (B–C) additive system has already been demonstrated to be effective for enhancing the densification progress [12]. In the conventional preparation process, the B sources are derived from boron powder or boron carbide, and the C sources are derived from carbon black or phenol resin. Usually, a slight excess of B and C additives is required to ensure the sintering of the SiC ceramics because the sintering aids are difficult to disperse uniformly [13]. However, due to the high covalent bond characteristics, SiC ceramics are difficult to densify without sintering aids [14]. In order to obtain dense and high-purity SiC ceramics, the boron and carbon content must be minimized in consideration of the solubility limit of B in SiC crystal lattice and the impurity O content of SiC green bodies [15].
Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai, 200050, China. Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai, 200050, China. E-mail addresses: [email protected] (Y. Yang), [email protected] (Z. Huang).
∗∗
https://doi.org/10.1016/j.ceramint.2019.06.231 Received 16 April 2019; Received in revised form 17 June 2019; Accepted 21 June 2019 Available online 26 June 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 19771–19776
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Fig. 1. Schematic illustration of preparation of the dense, high-purity SiC ceramics.
of 0.5 μm and a high purity of 99.86%. The main impurities contained 0.09 wt% free carbon, 0.03 wt% free silicon and 0.02 wt% metallic element (122 ppm Al, 64 ppm Ti and 14 ppm Fe). Boric acid (GR grade, Aladdin Chemistry Co., Ltd., Shanghai, China) with a B2O3 yield of about 55.34 wt% and D-fructose (BR grade, Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) with a C yield of about 17.81 wt% after thermolysis were used as B and C sources, respectively. 2.2. Process
Fig. 2. Effect of the ball milling time on O content of SiC green bodies of B4C18 after binder removal at 900 °C for 0.5 h.
The reported solubility of B in SiC grains is about 0.2 wt% at 2100 °C and less than 0.4 wt% at 2200 °C [16]. For S–SiC ceramics, the main source of impurities is residual B and C as the sintering additives due to their poor dispersibility. Therefore, in order to obtain high-purity SiC ceramics with high flexural strength for IC processing, ultra-low contents of additives with optimized mass fraction of B/C are used to prepare high-density SiC ceramics by pressureless solid-state-sintering, as shown in Fig. 1. In order to improve the dispersibility of the sintering additives, boric acid and D-fructose, which have higher solubility in ethanol, was doped into SiC powders as B–C sources. In this way, B4C, generated by in-situ reaction, can easily enter SiC lattices as a solid solution, which can better reduce the grain boundary energy. Besides, the added C as sintering aids should ensure that carbothermal reduction reaction occurs completely and make the residual C content minimum. In order to calculate that the amount of C needed as sintering additives in the carbothermal reduction reaction, the O content of SiC green compact after binder removal was measured. 2. Experimental procedure 2.1. Materials The starting powder was submicron α-SiC with a mean particle size
The SiC powders were mixed with the calculated amount of boric acid and D-fructose in ethanol to obtain the as-required recipe. The mixture was blended and wet ball milled for several hours at 300 r/min using SiC milling media. The slurry was dried using a rotary evaporator, then crushed and sieved with 100-mesh sieve. The as-sieved powders were uniaxially dry pressed in a steel die at 60 MPa, followed by cold isostatic pressing at 200 MPa. The green compacts were placed in a high-purity graphite crucible and heated to 900 °C under vacuum for binder removal. Finally, the samples were pressureless sintered in a graphite resistance furnace in argon. The temperature schedule was shown as the following: firstly heated to 1400 °C at 20 °C/min, then increased to1950 °C at 5 °C/min and kept for 0.5 h, and finally elevated to 2150 °C at 3 °C/min and maintained for 0.5 h. Fig. 2 shows the effect of the ball milling time on O content of SiC green bodies after binder removal. The amount of B and C as sintering additives was fixed on 0.4 wt% and 1.8 wt%, respectively. These samples were identified as B4C18. With the ball milling time increased, the O content of SiC green bodies was increased. In order to make the sintering aids fully dispersed homogeneously, the ball milling time was determined to be 24 h. When the ball milling was 24 h, the O content of B4C18 was 1.79 wt%, and mainly originated from B2O3 and SiO2. The residual B content was 0.16 wt% after binder removal measured by ICP-OES. Accordingly, the O content was 0.355 wt% from B2O3 and 1.435 wt% from SiO2, respectively. In the SiC green bodies, several major reactions which may occur are as follows:
2B2 O3 + 7C SiO2 + 3C
2SiO2 + SiC
B4 C + 6CO (g ) SiC + 2CO (g)
3SiO (g ) + CO (g )
(1) (2) (3)
According to the calculations by means of a commercial software (HSC Chemistry 6.0, Outokumpu Research Oy, Pori, Finland), reactions (1) and (2) could take place thermodynamically (ΔG < 0) in standard condition at temperatures higher than 1562.87 °C and 1521.38 °C, respectively. According to theoretical calculations, the amount of C needed was about 1.927 wt% when only reactions (1) and (2) were considered.
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Fig. 3. Gibbs free energy as a function of temperature corresponding to the carbothermal reduction reactions in the SiC green bodies.
2.3. Characterization Densities of green compacts and sintered ceramics were determined by the Archimedes method with kerosene and water as the immersing media, respectively. Densification degrees of as-sintered bulks were acquired by sintered densities over the theoretical density (3.21 g/cm3). Linear shrinkage in the whole sintering was calculated according to the length of a sample before and after sintered. The amount of free carbon was examined by Carbon and Sulfur Analyzer (CS-2000; Eltra, Germany) and the O content was measured by Oxygen and Nitrogen Analyzer (TC600C; Leco, USA). In addition, the B content was examined by Inductively Coupled Plasma-Optical Emission Spectrometer (ICPOES; 725 ICP-OES; NYSE: A, USA). The phase composition was determined by standard powder X-ray diffraction (XRD; D8 ADVANCE; BRUKER, Germany) using Cu Ka radiation. The microstructures of ceramics after polished and etched were observed with field emission scanning electron microscopy (SEM; JSM-6700F; JEOL, Japan). Finally, The purity was analysed by Glow Discharge Mass Spectrometry (GDMS; AUTOCONCEPT-GD90RF; MSI, UK). 3. Results and discussion 3.1. Effect of added C content Fig. 4 shows the effect of the amount of C as sintering additives on densification degree and the residual C content of as-sintered S–SiC ceramics. The amount of B as sintering additives was fixed on 0.4 wt%. As the added C content increased, the relative density firstly increased and then did almost not changed. When the content of added C was less than 1.8 wt%, the B2O3 and SiO2 can not be reacted completely. For one thing, there was not enough B4C to reduce the grain boundary energy. For another thing, because a portion of the SiO2 remained on the surface of SiC powders, the surface energy was not high enough for sintering that the green bodies were unable to be dense. These samples, whose added C content was 1.8 wt%, had the lowest residual C content with 0.14 wt%. On the one hand, when the added C content was less than 1.8 wt%, all the added C had been consumed in the carbothermal reduction reaction. And the graphite vapor, which was coming from graphite furnace and graphite crucible, would be adsorbed into the pores largely existed in those samples (Fig. 5a and b). Therefore, the residual C content decreased with the relative density increased. On the other hand, when the added C content was more than 1.8 wt%, after the reduction of B2O3 and SiO2, there still was graphite
Fig. 4. Effect of the amount of C as sintering additives on densification degree and the residual C content of as-sintered S–SiC ceramics at 2150 °C for 0.5 h.
remained as a secondary phase at the grain boundaries (Fig. 5c). Compared with the theoretical value, the experiment value of C needed as sintering additives was slightly less because of reaction (3). 3.2. Effect of added B content Fig. 6 shows the effect of the amount of B as sintering additives on densification degree and the residual B and C content of as-sintered S–SiC ceramics. The amount of C as sintering aids was fixed on 1.8 wt%. With the added B content increased, the densification degree firstly increased and then decreased; exactly the opposite, the residual C content firstly decreased and then increased. When the added B content reached 0.4 wt%, the relative density peaked up to 98.42% and the residual C content reached the nadir with 0.14 wt%. For other samples, there was a common ground that the graphite vapor would be adsorbed into the pores due to their lower density, which would make the residual C content higher (Fig. 7a–e). Furthermore, these samples, which the added B content was less than 0.4 wt%, had another reason for higher residual C content that the added C content was more than the amount of C needed in the carbothermal reduction reaction. The reason, why the relative density decreased when the added B content was higher than 0.4 wt%, was that the added C was not enough to redox all of the B2O3 and SiO2. Fig. 3 shows the Gibbs free energy of reactions (1) and (2). The initial reaction temperatures of the two reactions are not much different, and there is no obvious priority for the two reactions. This results in a small amount of residual SiO2 preventing the densification of the samples. The residual B content almost increased as the added B content increased. The reason why residual B content was not equal to B content, was because of the volatility of B2O3. The residual B content of these samples was only 0.16 wt% after binder removal and 0.14 wt% after sintering, which was 0.4 wt% in initial samples. 3.3. The microstructures and properties of B4C18 samples Fig. 8 shows the SEM and BSD images of the B4C18 samples. The grains were all columnar with high aspect ratio (2:1–10:1), and the length of some of them was more than 450 μm. Because there was no large second phase such as free carbon or B4C at grain boundaries and no pinning effect that came from grain grading, to suppress the abnormal grain growth in these samples [11]. Besides, there were a lot of black dots in the gigantic grains, and most of them were tiny graphite particles. Because D-fructose had excellent dispersibility, and the
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Fig. 5. Microstructures of SiC ceramics with 0.4 wt% B and different content of C: (a) 1.2 wt%; (b) 1.5 wt%; (c) 2.4 wt%.
Fig. 6. Effect of the amount of B as sintering additives on densification degree and the residual B and C content of as-sintered S–SiC ceramics at 2150 °C for 0.5 h.
graphite particles was too tiny to suppress the grain growth, so that they were surrounded into SiC grains. XRD result (Fig. 9) showed a hexagonal crystal structure of 4H–SiC
(PDF #73–1664), where the obvious diffraction peaks at 33.7°(100), 34.9°(101), 35.8°(004), 38.3°(102), 43.4°(103), 49.8°(104), 60.1°(110), 71.9°(201), 73.5°(114) and 73.7°(202), and 6H–SiC (PDF #75–1541), where the obvious diffraction peaks at 34.3°(101), 35.8°(006), 38.3°(103), 41.5°(104), 45.5°(105), 60.1°(108), 73.5°(116) and 73.7°(203), in the specimens of B4C18. These was no characteristic diffraction peak of graphite in the picture. In addition, chemical analysis was conducted by GDMS. Table 1 shows elementary analysis of the samples of B4C18, and only elements with a content greater than 0.1 ppm are listed in this table. Except B and O, the total amount of other impurities was 145.08 ppm. The purity of the as-prepared SiC ceramics were about 99.6 wt%. The green density of B4C18 specimens is 1.798 g/cm3, and the linear shrinkage is 23.3%. The flexural strength of the dense, highpurity SiC Ceramics is 443 ± 27 MPa, and the elasticity modulus of them was 420 ± 1 GPa. Furthermore, the coefficient of thermal expansion of them at RT to 600 °C is 3.84 × 10−6 K−1. 4. Conclusions In summary, the effect of different contents of sintering additives on SiC ceramics has been investigated, and dense, high-purity SiC ceramics have been obtained by pressureless sintering. In the possess, the amount of boric acid and D-fructose was calculated to generate 0.4 wt% B and
Fig. 7. Microstructures of SiC ceramics with 1.8 wt% C and different content of B: (a) 0.3 wt%; (b) 0.35 wt%; (c) 0.45 wt%; (d) 0.5 wt%; (e) 0.6 wt%.
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Fig. 8. Microstructures of the B4C18 specimens: (a) the SEM image after polished and corroded by NaOH; (b), (c) and (d) the BSD image after polished and etched. Table 1 Elementary analysis of as-sintered S–SiC ceramics of B4C18 at 2150 °C for 0.5 h.
Fig. 9. XRD patterns of as-sintered S–SiC ceramics of B4C18 at 2150 °C for 0.5 h.
1.8 wt% C as sintering aids by carbothermal reduction reaction, and the residual B and C content were 0.114 wt% and 0.14 wt%, respectively. Besides, the amount of O in the as-prepared SiC ceramics was 0.136 wt %, and the total content of other impurity elements (such as aluminum, magnesium, calcium, iron, etc.) was 145.08 ppm. Finally, the purity of the as-prepared SiC ceramics was more than 99.5 wt%.
Element
Concentration (ppm)
B C O Mg Al Si P S K Ca Ti Cr Mn Fe Ni Zn
1136 Matrix 1365 11 89 Matrix 7.9 0.48 < 0.5 18 0.98 1.3 0.13 15 0.48 0.31
Acknowledgment The authors gratefully acknowledge the finical support of the National Key Research and Development Project (No. 2017YFB0310600), and this work is also supported by Shanghai international science and Technology Cooperation Fund (No. 17520711700 and 18520744200).
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References [1] E. Ciudad, O. Borrero-López, F. Rodríguez-Rojas, et al., Effect of intergranular phase chemistry on the sliding-wear resistance of pressureless liquid-phase-sintered α-SiC, J. Eur. Ceram. Soc. 32 (2) (2012) 511–516. [2] L. Charpentier, A. Maître, M. Balat-Pichelin, et al., Influence of alumina on the passive oxidation at low oxygen pressure of hot-pressed α-SiC, Scripta Mater. 60 (7) (2009) 481–484. [3] F. Rodríguez-Rojas, A.L. Ortiz, F. Guiberteau, et al., Anomalous oxidation behaviour of pressureless liquid-phase-sintered SiC, J. Eur. Ceram. Soc. 31 (13) (2011) 2393–2400. [4] F. Rodríguez-Rojas, A.L. Ortiz, F. Guiberteau, et al., Oxidation behaviour of pressureless liquid-phase-sintered α-SiC with additions of 5Al2O3 + 3RE2O3 (RE = La, Nd, Y, Er, Tm, or Yb), J. Eur. Ceram. Soc. 30 (15) (2010) 3209–3217. [5] A.L. Ortiz, O. Borrero-López, M.Z. Quadir, et al., A route for the pressureless liquidphase sintering of SiC with low additive content for improved sliding-wear resistance, J. Eur. Ceram. Soc. 32 (4) (2012) 965–973. [6] W. Guo, H. Xiao, W. Xie, et al., A new design for preparation of high performance recrystallized silicon carbide, Ceram. Int. 38 (3) (2012) 2475–2481. [7] S. Chiang, An application of Lotka–Volterra model to Taiwan's transition from 200 mm to 300 mm silicon wafers, Technol. Forecast. Soc. Change 79 (2) (2012) 383–392.
[8] Z. Lu, S. Kimbel, Growth of 450 mm diameter semiconductor grade silicon crystals, Journal of Crystal Growth 318 (2011) 193–195. [9] J.S. Nadeau, Very high pressure hot pressing of silicon carbide, Am. Ceram. Soc. Bull. 52 (2) (1973) 170–174. [10] A. Lara, A.L. Ortiz, A. Munoz, A. Dominguez-Rodriguez, Densification of additivefree polycrystalline β-SiC by spark-plasma sintering, Ceram. Int. 38 (1) (2012) 45–53. [11] Haibo Wu, Yongjie Yan, Guilin Liu, et al., Effects of grain grading on microstructures and mechanical behaviors of pressureless solid-state-sintered SiC, Int. J. Appl. Ceram. Technol. 12 (5) (2015) 976–984. [12] E.R. Maddrell, Pressureless sintering of silicon carbide, J. Mater. Sci. Lett. 6 (4) (1987) 486–488. [13] W.J. Clegg, Role of carbon in the sintering of boron-doped silicon carbide, J. Am. Ceram. Soc. 83 (5) (2000) 1039–1043. [14] E.J. Lee, H.L. Dong, J.C. Kim, et al., Densification behavior of high purity SiC by hot pressing, Ceram. Int. 40 (10) (2014) 16389–16392. [15] Y. Li, J. Yin, H. Wu, et al., High thermal conductivity in pressureless densified SiC ceramics with ultra-low contents of additives derived from novel boron-carbon sources, J. Eur. Ceram. Soc. 34 (10) (2014) 2591–2595. [16] Y. Murata, R.H. Smoak, Factors in densification and sintering of oxide and nonoxide ceramics, Proceedings of international symposium (1978) 382–399.
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Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Preparation of dense and high-purity SiC ceramics by pressureless solidstate-sintering
T
Meng Liua,b, Yong Yanga,c,∗, Yuquan Weia, Yajie Lia,b, Heng Zhanga, Xuejian Liua,c, Zhengren Huanga,b,∗∗ a
Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China School of Physical Science and Technology, ShanghaiTech University, Shanghai, 201210, China c Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing, 100049, China b
ARTICLE INFO
ABSTRACT
Keywords: SiC Dense High-purity Pressureless solid-state-sintering
High-purity SiC materials have been used in semiconductor processes due to their excellent properties. However, they are difficult to densify without sintering aids. In this work, dense and high-purity SiC ceramics have been obtained by pressureless solid-state-sintering with ultra-low contents of sintering additives. The amount of residual B, C and O in the high-purity SiC ceramics was less than 0.15 wt%, respectively, and the total content of other impurity elements (such as aluminum, magnesium, calcium, iron, etc.) was less than 0.015 wt%. Finally, the purity of the as-prepared SiC ceramics was more than 99.5 wt%.
1. Introduction Silicon carbide (SiC) ceramics, due to strong covalent Si–C bond, possess many excellent properties such as high strength, good hardness, excellent chemical stability, outstanding oxidation and wear resistance [1–5]. Therefore, high-purity SiC ceramic devices are widely used in the various processing furnaces for silicon wafers (single crystal growth, oxidation, corrosion, ion injection, low-pressure CVD, etc.) [6]. Presently, silicon wafers with the diameter of 300 nm have been massively applied in IC manufacturing, and the wafers of 450 mm will be introduced in the near future due to the effect of Moore's Law. The larger wafers must be thicker to increase their resistance to warping and other structural deformities [7,8]. Therefore, there is a higher requirement of high-purity SiC ceramic devices with large size and high mechanical property used in the silicon wafer processing. Usually, large sized and complex shaped high-purity SiC ceramic devices can be fabricated by recrystallization sintering. Nevertheless, recrystallized silicon carbide (R–SiC) has lower flexural strength and poor oxidation resistance because of its relatively low density and porous structure. For example, Wenming Guo et al. [6] fabricated a high performance R–SiC with a density of 2.99 g/cm3 and flexural strength of 162.3 ± 2.6 MPa which was obtained by three polymer impregnation and pyrolysis-recrystallization cycles. But even so, the flexural strength of R–SiC can still hardly meet the requirement of the semiconductor industry.
∗
Many investigations have been done to increase the density of SiC ceramics without the sintering additive. Nadeau et al. [9] and Lara et al. [10] fabricated fully dense SiC bodies without sintering aids by hotpressed sintering at 2500 °C, 5000 MPa and spark plasma sintering at 2100 °C, 70 MPa, respectively. Nevertheless, neither of these methods is suitable for preparing SiC ceramic devices of large size and complex shapes. Pressureless solid-state-sintered SiC (S–SiC) could be dense and high-purity, and can be easily processed into ceramic devices with large size and complex shapes. Especially, it possesses a unique combination of properties including high thermal conductivity and outstanding hightemperature strength [11]. For S–SiC ceramics, boron-carbon (B–C) additive system has already been demonstrated to be effective for enhancing the densification progress [12]. In the conventional preparation process, the B sources are derived from boron powder or boron carbide, and the C sources are derived from carbon black or phenol resin. Usually, a slight excess of B and C additives is required to ensure the sintering of the SiC ceramics because the sintering aids are difficult to disperse uniformly [13]. However, due to the high covalent bond characteristics, SiC ceramics are difficult to densify without sintering aids [14]. In order to obtain dense and high-purity SiC ceramics, the boron and carbon content must be minimized in consideration of the solubility limit of B in SiC crystal lattice and the impurity O content of SiC green bodies [15].
Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai, 200050, China. Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai, 200050, China. E-mail addresses: [email protected] (Y. Yang), [email protected] (Z. Huang).
∗∗
https://doi.org/10.1016/j.ceramint.2019.06.231 Received 16 April 2019; Received in revised form 17 June 2019; Accepted 21 June 2019 Available online 26 June 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Ceramics International 45 (2019) 19771–19776
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Fig. 1. Schematic illustration of preparation of the dense, high-purity SiC ceramics.
of 0.5 μm and a high purity of 99.86%. The main impurities contained 0.09 wt% free carbon, 0.03 wt% free silicon and 0.02 wt% metallic element (122 ppm Al, 64 ppm Ti and 14 ppm Fe). Boric acid (GR grade, Aladdin Chemistry Co., Ltd., Shanghai, China) with a B2O3 yield of about 55.34 wt% and D-fructose (BR grade, Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) with a C yield of about 17.81 wt% after thermolysis were used as B and C sources, respectively. 2.2. Process
Fig. 2. Effect of the ball milling time on O content of SiC green bodies of B4C18 after binder removal at 900 °C for 0.5 h.
The reported solubility of B in SiC grains is about 0.2 wt% at 2100 °C and less than 0.4 wt% at 2200 °C [16]. For S–SiC ceramics, the main source of impurities is residual B and C as the sintering additives due to their poor dispersibility. Therefore, in order to obtain high-purity SiC ceramics with high flexural strength for IC processing, ultra-low contents of additives with optimized mass fraction of B/C are used to prepare high-density SiC ceramics by pressureless solid-state-sintering, as shown in Fig. 1. In order to improve the dispersibility of the sintering additives, boric acid and D-fructose, which have higher solubility in ethanol, was doped into SiC powders as B–C sources. In this way, B4C, generated by in-situ reaction, can easily enter SiC lattices as a solid solution, which can better reduce the grain boundary energy. Besides, the added C as sintering aids should ensure that carbothermal reduction reaction occurs completely and make the residual C content minimum. In order to calculate that the amount of C needed as sintering additives in the carbothermal reduction reaction, the O content of SiC green compact after binder removal was measured. 2. Experimental procedure 2.1. Materials The starting powder was submicron α-SiC with a mean particle size
The SiC powders were mixed with the calculated amount of boric acid and D-fructose in ethanol to obtain the as-required recipe. The mixture was blended and wet ball milled for several hours at 300 r/min using SiC milling media. The slurry was dried using a rotary evaporator, then crushed and sieved with 100-mesh sieve. The as-sieved powders were uniaxially dry pressed in a steel die at 60 MPa, followed by cold isostatic pressing at 200 MPa. The green compacts were placed in a high-purity graphite crucible and heated to 900 °C under vacuum for binder removal. Finally, the samples were pressureless sintered in a graphite resistance furnace in argon. The temperature schedule was shown as the following: firstly heated to 1400 °C at 20 °C/min, then increased to1950 °C at 5 °C/min and kept for 0.5 h, and finally elevated to 2150 °C at 3 °C/min and maintained for 0.5 h. Fig. 2 shows the effect of the ball milling time on O content of SiC green bodies after binder removal. The amount of B and C as sintering additives was fixed on 0.4 wt% and 1.8 wt%, respectively. These samples were identified as B4C18. With the ball milling time increased, the O content of SiC green bodies was increased. In order to make the sintering aids fully dispersed homogeneously, the ball milling time was determined to be 24 h. When the ball milling was 24 h, the O content of B4C18 was 1.79 wt%, and mainly originated from B2O3 and SiO2. The residual B content was 0.16 wt% after binder removal measured by ICP-OES. Accordingly, the O content was 0.355 wt% from B2O3 and 1.435 wt% from SiO2, respectively. In the SiC green bodies, several major reactions which may occur are as follows:
2B2 O3 + 7C SiO2 + 3C
2SiO2 + SiC
B4 C + 6CO (g ) SiC + 2CO (g)
3SiO (g ) + CO (g )
(1) (2) (3)
According to the calculations by means of a commercial software (HSC Chemistry 6.0, Outokumpu Research Oy, Pori, Finland), reactions (1) and (2) could take place thermodynamically (ΔG < 0) in standard condition at temperatures higher than 1562.87 °C and 1521.38 °C, respectively. According to theoretical calculations, the amount of C needed was about 1.927 wt% when only reactions (1) and (2) were considered.
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Fig. 3. Gibbs free energy as a function of temperature corresponding to the carbothermal reduction reactions in the SiC green bodies.
2.3. Characterization Densities of green compacts and sintered ceramics were determined by the Archimedes method with kerosene and water as the immersing media, respectively. Densification degrees of as-sintered bulks were acquired by sintered densities over the theoretical density (3.21 g/cm3). Linear shrinkage in the whole sintering was calculated according to the length of a sample before and after sintered. The amount of free carbon was examined by Carbon and Sulfur Analyzer (CS-2000; Eltra, Germany) and the O content was measured by Oxygen and Nitrogen Analyzer (TC600C; Leco, USA). In addition, the B content was examined by Inductively Coupled Plasma-Optical Emission Spectrometer (ICPOES; 725 ICP-OES; NYSE: A, USA). The phase composition was determined by standard powder X-ray diffraction (XRD; D8 ADVANCE; BRUKER, Germany) using Cu Ka radiation. The microstructures of ceramics after polished and etched were observed with field emission scanning electron microscopy (SEM; JSM-6700F; JEOL, Japan). Finally, The purity was analysed by Glow Discharge Mass Spectrometry (GDMS; AUTOCONCEPT-GD90RF; MSI, UK). 3. Results and discussion 3.1. Effect of added C content Fig. 4 shows the effect of the amount of C as sintering additives on densification degree and the residual C content of as-sintered S–SiC ceramics. The amount of B as sintering additives was fixed on 0.4 wt%. As the added C content increased, the relative density firstly increased and then did almost not changed. When the content of added C was less than 1.8 wt%, the B2O3 and SiO2 can not be reacted completely. For one thing, there was not enough B4C to reduce the grain boundary energy. For another thing, because a portion of the SiO2 remained on the surface of SiC powders, the surface energy was not high enough for sintering that the green bodies were unable to be dense. These samples, whose added C content was 1.8 wt%, had the lowest residual C content with 0.14 wt%. On the one hand, when the added C content was less than 1.8 wt%, all the added C had been consumed in the carbothermal reduction reaction. And the graphite vapor, which was coming from graphite furnace and graphite crucible, would be adsorbed into the pores largely existed in those samples (Fig. 5a and b). Therefore, the residual C content decreased with the relative density increased. On the other hand, when the added C content was more than 1.8 wt%, after the reduction of B2O3 and SiO2, there still was graphite
Fig. 4. Effect of the amount of C as sintering additives on densification degree and the residual C content of as-sintered S–SiC ceramics at 2150 °C for 0.5 h.
remained as a secondary phase at the grain boundaries (Fig. 5c). Compared with the theoretical value, the experiment value of C needed as sintering additives was slightly less because of reaction (3). 3.2. Effect of added B content Fig. 6 shows the effect of the amount of B as sintering additives on densification degree and the residual B and C content of as-sintered S–SiC ceramics. The amount of C as sintering aids was fixed on 1.8 wt%. With the added B content increased, the densification degree firstly increased and then decreased; exactly the opposite, the residual C content firstly decreased and then increased. When the added B content reached 0.4 wt%, the relative density peaked up to 98.42% and the residual C content reached the nadir with 0.14 wt%. For other samples, there was a common ground that the graphite vapor would be adsorbed into the pores due to their lower density, which would make the residual C content higher (Fig. 7a–e). Furthermore, these samples, which the added B content was less than 0.4 wt%, had another reason for higher residual C content that the added C content was more than the amount of C needed in the carbothermal reduction reaction. The reason, why the relative density decreased when the added B content was higher than 0.4 wt%, was that the added C was not enough to redox all of the B2O3 and SiO2. Fig. 3 shows the Gibbs free energy of reactions (1) and (2). The initial reaction temperatures of the two reactions are not much different, and there is no obvious priority for the two reactions. This results in a small amount of residual SiO2 preventing the densification of the samples. The residual B content almost increased as the added B content increased. The reason why residual B content was not equal to B content, was because of the volatility of B2O3. The residual B content of these samples was only 0.16 wt% after binder removal and 0.14 wt% after sintering, which was 0.4 wt% in initial samples. 3.3. The microstructures and properties of B4C18 samples Fig. 8 shows the SEM and BSD images of the B4C18 samples. The grains were all columnar with high aspect ratio (2:1–10:1), and the length of some of them was more than 450 μm. Because there was no large second phase such as free carbon or B4C at grain boundaries and no pinning effect that came from grain grading, to suppress the abnormal grain growth in these samples [11]. Besides, there were a lot of black dots in the gigantic grains, and most of them were tiny graphite particles. Because D-fructose had excellent dispersibility, and the
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Fig. 5. Microstructures of SiC ceramics with 0.4 wt% B and different content of C: (a) 1.2 wt%; (b) 1.5 wt%; (c) 2.4 wt%.
Fig. 6. Effect of the amount of B as sintering additives on densification degree and the residual B and C content of as-sintered S–SiC ceramics at 2150 °C for 0.5 h.
graphite particles was too tiny to suppress the grain growth, so that they were surrounded into SiC grains. XRD result (Fig. 9) showed a hexagonal crystal structure of 4H–SiC
(PDF #73–1664), where the obvious diffraction peaks at 33.7°(100), 34.9°(101), 35.8°(004), 38.3°(102), 43.4°(103), 49.8°(104), 60.1°(110), 71.9°(201), 73.5°(114) and 73.7°(202), and 6H–SiC (PDF #75–1541), where the obvious diffraction peaks at 34.3°(101), 35.8°(006), 38.3°(103), 41.5°(104), 45.5°(105), 60.1°(108), 73.5°(116) and 73.7°(203), in the specimens of B4C18. These was no characteristic diffraction peak of graphite in the picture. In addition, chemical analysis was conducted by GDMS. Table 1 shows elementary analysis of the samples of B4C18, and only elements with a content greater than 0.1 ppm are listed in this table. Except B and O, the total amount of other impurities was 145.08 ppm. The purity of the as-prepared SiC ceramics were about 99.6 wt%. The green density of B4C18 specimens is 1.798 g/cm3, and the linear shrinkage is 23.3%. The flexural strength of the dense, highpurity SiC Ceramics is 443 ± 27 MPa, and the elasticity modulus of them was 420 ± 1 GPa. Furthermore, the coefficient of thermal expansion of them at RT to 600 °C is 3.84 × 10−6 K−1. 4. Conclusions In summary, the effect of different contents of sintering additives on SiC ceramics has been investigated, and dense, high-purity SiC ceramics have been obtained by pressureless sintering. In the possess, the amount of boric acid and D-fructose was calculated to generate 0.4 wt% B and
Fig. 7. Microstructures of SiC ceramics with 1.8 wt% C and different content of B: (a) 0.3 wt%; (b) 0.35 wt%; (c) 0.45 wt%; (d) 0.5 wt%; (e) 0.6 wt%.
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Fig. 8. Microstructures of the B4C18 specimens: (a) the SEM image after polished and corroded by NaOH; (b), (c) and (d) the BSD image after polished and etched. Table 1 Elementary analysis of as-sintered S–SiC ceramics of B4C18 at 2150 °C for 0.5 h.
Fig. 9. XRD patterns of as-sintered S–SiC ceramics of B4C18 at 2150 °C for 0.5 h.
1.8 wt% C as sintering aids by carbothermal reduction reaction, and the residual B and C content were 0.114 wt% and 0.14 wt%, respectively. Besides, the amount of O in the as-prepared SiC ceramics was 0.136 wt %, and the total content of other impurity elements (such as aluminum, magnesium, calcium, iron, etc.) was 145.08 ppm. Finally, the purity of the as-prepared SiC ceramics was more than 99.5 wt%.
Element
Concentration (ppm)
B C O Mg Al Si P S K Ca Ti Cr Mn Fe Ni Zn
1136 Matrix 1365 11 89 Matrix 7.9 0.48 < 0.5 18 0.98 1.3 0.13 15 0.48 0.31
Acknowledgment The authors gratefully acknowledge the finical support of the National Key Research and Development Project (No. 2017YFB0310600), and this work is also supported by Shanghai international science and Technology Cooperation Fund (No. 17520711700 and 18520744200).
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