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Microstructures and properties of SiB0.5C1.5N0.5 ceramics consolidated by mechanical alloying and hot pressing
Materials Science and Engineering A 489 (2008) 187–192
Microstructures and properties of SiB0.5C1.5N0.5 ceramics consolidated by mechanical alloying and hot pressing Zhi-Hua Yang ∗ , Yu Zhou, De-Chang Jia, Qing-Chang Meng Institute for Advanced Ceramics, Harbin Institute of Technology, Harbin 150080, PR China Received 27 August 2007; received in revised form 7 December 2007; accepted 7 December 2007
Abstract Nanostructured SiB0.5 C1.5 N0.5 ceramics were formed by two routes of mechanical alloying and consolidated by hot pressing. Nano-sized 6H-SiC, 3C-SiC and BCN were the main phases for two SiB0.5 C1.5 N0.5 ceramics. For the SiBCN ceramics fabricated from the powders milled by one-step milling, the content of 6H-SiC is higher than that of ceramic fabricated from the powders milled by two-step milling. The flexural strengths for the former and latter were 312.8 MPa and 423.4 MPa at room temperature, respectively. At high temperature, the strengths of the former reduced 8.2% and 28.0% at 1000 ◦ C and 1400 ◦ C, respectively. The reduction rates of strengths for the latter were 11.3% and 38.5%, respectively. The average thermal expansion coefficients of two SiBCN ceramics consolidated at 1900 ◦ C were about 4.5 × 10−6 /◦ C. © 2008 Elsevier B.V. All rights reserved. Keywords: SiB0.5 C1.5 N0.5 ceramics; Mechanical alloying; Hot pressing; Mechanical properties; Microstructure
1. Introduction SiBCN ceramics derived from polymer precursors recently received a significant attention because of their outstanding hightemperature stability [1] and mechanical properties [2]. SiBCN ceramics had low density and the lowest reported oxidation rates of any non-oxide material known to date [3], so those ceramics suggest themselves for applications under extreme conditions like in heat engines. Some works have been done to develop such materials from the stage of laboratory curiosities to engineering materials by attempting industrial-scale processing, and, hence, it is very important to know the mechanical and thermal properties of SiBCN ceramics. However, up to now, only hightemperature compression creep properties were reported [4–7]. It may be due to the method of polymer precursors limiting the fabrications of big size bulk ceramics. Powder processing by mechanical alloying (MA) has attracted wide practical interest as it offers a simple but powerful way to synthesize non-equilibrium phases and microstructures, from nanograin materials to extended solid solutions, amorphous phases, chemically disordered compounds, and nanocomposites [8]. SiC [9–12], Si3 N4 –SiC [13], and BCN ∗
Corresponding author. Tel.: +86 451 86418792; fax: +86 451 86414291. E-mail address: [email protected] (Z.-H. Yang).
0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.12.010
[14–16] powders have been fabricated by MA. Our previous work established that nanostructured composite SiB0.5 C1.5 N0.5 can be produced by MA [17]. In the present paper, SiBCN ceramics were consolidated by hot pressing (HP) using two types of milled SiBCN powders as starting materials. Microstructures of SiBCN ceramics were characterized. Mechanical properties (at room and high temperature) and thermal expansion coefficient were also tested. 2. Experimental procedure SiB0.5 C1.5 N0.5 powders (with a target composition of SiC–33.3 at.% BCN) were fabricated by one- and two-step MA using shake mill [17,18]. For one-step milling, c-Si, h-BN and graphite (C) powders were simultaneously milled for 20 h. As for two-step milling, both Si and partial C powders in mole ratio of 1:1 were milled for 15 h, and then the residual C and BN powders were added into the vessel for another 5 h milling. The powders prepared by one- and two-step milling were denoted as CS and CD, respectively. Two SiB0.5 C1.5 N0.5 powders mainly contained amorphous SiBCN, nano-crystalline (4–5 nm) 3CSiC and 6H-SiC. In addition, about 4.7% Si remained in the CS powders. Sintering was performed by HP at 1850/1900 ◦ C under a pressure of 40 MPa for 30 min in N2 (1.0 × 105 Pa). The furnace for
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Fig. 1. XRD patterns of the bulk samples produced by HP of SiBCN powders at various holding temperatures.
tent with the analysis results of the milled SiBCN powders [18]. In addition, the appearance of ZrO2 reflections was due to the contamination from the wear of ZrO2 balls and inner surface of vessel. The relative density is 91.5% for the CS1850, and it can reach 99.4% for CS1900. However, the relative density of CD1900 is only 94.3%. SEM photographs of polishing surface of SiBCN ceramics are shown in Fig. 2. As can be seen that the size of all grains are <1 m. Fine, equiaxed SiC grains (light phase) were homogeneously distributed in SiBCN ceramics. Many micro cracks were found the around the SiC grains, and some cracks even crossed the SiC grains, as shown by the white circle in the Fig. 2c. Micro cracks may be caused by the thermal shock during the module transferred from the sintering chamber to depositing chamber. The TEM images of CS and CD powders consolidated at different temperatures are shown in Fig. 3. For the CS1850, SiC
the sintering has two chambers. One chamber is for the sintering, and the other one is used to deposit module which was transferred from the sintering chamber. The temperature of depositing chamber is below 100 ◦ C during the sintering. The module was transferred into the depositing chamber when the sintering temperature decreased to 900 ◦ C in order to shorten the duration of sintering cycle. Hereafter, the sintered samples were marked as CS/CD#, # = 1850 or 1900, corresponding to consolidation temperatures of 1850 and 1900 ◦ C. The resulting sample ceramics were cut into bars of 36 mm × 3 mm × 4 mm (30 mm outer span) for measuring flexural strength at room and high temperature with a crosshead speed of 0.5 mm/min. The samples were heated up to 1000 and 1400 ◦ C at a rate of 10 ◦ C/min in air, and held at chosen temperature for 10 min before loading to ensure thermal equilibrium. The fracture toughness was determined using the single edge notched beam method with a crosshead speed of 0.05 mm/min. The testing bar dimensions used were 2 mm × 4 mm × 20 mm (16 mm outer span). The depth of the notches was 2.0 mm and the width about 0.2 mm. The structural characterization of ceramics was analyzed using X-ray diffraction (XRD) methods with Cu K␣ radiation. The investigation of microstructure was done in scanning electron microscope (SEM) and transmission electron microscope (TEM) operating at 120 kV. Thermogravimetric analysis (TGA) of the ceramics was performed in Al2 O3 crucible (STA-449C, Netzsch) by heating up to 1500 ◦ C at a heating rate of 10 ◦ C/min in air. Thermal expansion coefficient was performed by heating up to 1400 ◦ C at a heating rate of 10 ◦ C/min in argon. 3. Results and discussion Fig. 1 shows the XRD patterns of the bulk samples produced by HP of SiBCN powders at various holding temperatures. In all cases, clear diffraction peaks are observed suggesting that most grains have recrystallized during the sintering process. BCN, 3CSiC and 6H-SiC are the main phase for the CS and CD ceramics. 3C-SiC is the main SiC phase in CD1900, but 6H-SiC is predominant phase for CS ceramics. The coexistence of 3C/6H-SiC and their relative contents for the two SiBCN ceramics were consis-
Fig. 2. SEM images of polishing surface of SiBCN ceramics hot-pressed at different temperature. (a) CS1850, (b) CS1900, and (c) CD1900.
Z.-H. Yang et al. / Materials Science and Engineering A 489 (2008) 187–192
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Fig. 3. TEM images of CS and CD powders sintered at different temperatures. (a) CS1850, (b) CS1900, and (c) CD1900.
grain sizes range from 100 nm to 300 nm (as shown in Fig. 3a). BCN mainly distribute the grains boundaries of SiC crystal. The grain sizes of BCN phase are very small, and it is difficult to display the BCN nano-crystal by the TEM. Grain sizes of SiC crystal increased up to 500 nm when the sintering temperature increasing 50 ◦ C. And the Selected Area Electron Diffraction (SAD) patterns from the selected area b-1 and b-2 exhibit [0 0 1] and [0 1 1¯ 0] patterns indicative of the -3C and ␣-6H structures. Some of BCN crystal visibly formed, and grain sizes increased for the CS1900 (as shown in Fig. 3b). For the CD1900, the grain sizes of -SiC are similar with these of CS1900 (as shown in Fig. 3c). BCN phase do not have visible crystals, but the size of BCN area is bigger than that in CS1850 and CS1900, which may be caused by the shorter times of milling (5 h) for BN-C powders in CD than these (20 h) in CS. The surface chemistry of CS1900 platelet was characterized by XPS. The measurement results are shown in Fig. 4. The characteristic peaks at 99.8 eV and 284.3 eV were identified as Si2p and C1s signals of the SiC covalent bond, respectively [19]. The
peaks at 191.0 eV (B1s), 285.2 eV (C1s), and 400.8 eV (N1s) can reflect the existence of B–C–N bond [20]. The peaks at 190.6 eV (B1s) and 397.7 eV (N1s) may be assigned to the B–N bonds of h-BN [20]. There is a peak at 284.3 eV for C1s spectrum, which assigned to the C–C bonds of graphite and agreed well with result of starting graphite [21]. The peak at 101.4 eV for Si2p spectrum was assigned to the Si–N bonds [19], but the content of Si–N bonds was <4% according the calculating method of SiBCN powders [18]. Therefore, it is difficult to detect by the XRD for the Si–N bonds. Structural characterization by XPS recorded the coexistence of several bonds between Si, B, C and N atoms. The synthesized SiBCN ceramics appeared to possess a structure containing several bonds such as Si–C, Si–N, B–C–N, C–C with B–N hybridization, rather than a simple complex of these bonds. Mechanical properties of SiBCN ceramics at room and high temperature are shown in Table 1. It can be seen that the sintering temperature and content of ␣-SiC affected mechanical properties of SiBCN ceramics. Mechanical properties of CS
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Fig. 4. (a) Si2p, (b) B1s, (c) C1s and (d) N1s XPS spectra, with deconvolutions of peaks, recorded for CS1900 ceramics. Table 1 Mechanical properties of SiBCN ceramics Sample
CS1850 CS1900 CD1900
Flexural strength (MPa) RT
1000 ◦ C
1400 ◦ C
191.7 ± 20.7 312.8 ± 5.2 423.4 ± 43.0
– 287.2 ± 15.5 375.4 ± 10.7
– 225.3 ± 15.1 260.4 ± 15.7
Elastic modulus (GPa)
Fracture toughness (MPa m1/2 )
Vickers hardness (GPa)
107.9 ± 9.5 136.3 ± 17.8 134.6 ± 13.5
1.80 ± 0.07 3.31 ± 0.02 3.09 ± 0.05
2.44 ± 0.31 4.17 ± 0.50 3.89 ± 0.43
ceramics increased with the increase of sintering temperature. The increase of properties could be contributed to the increase of relative density from 91.5% to 99.4%. As for the CS1900 and CD1900, the content of ␣-SiC have a great influence on the mechanical properties. The microstructure of SiC obtained from an ␣-SiC powder is composed of fine equiaxed grains and is usually brittle. And the ceramics, which have higher content of ␣-SiC, have higher fracture toughness [22]. So it can be seen that the flexural strength of CD1900 is 110 MPa higher than that of CS1900 at room temperature (RT), despite the low density of CD1900. The fracture toughness of CS1900 is higher than that of CD1900. CD1900 exhibits higher strengths at RT and high temperature than CS1900. High-temperature strengths of CD1900 reduced 11.3% at 1000 ◦ C and 38.5% at 1400 ◦ C, respectively, but the values were only 8.2% and 28.0% for CS1900. The reasons for the lower reduction rates of high-temperature strength of CS1900 may be related to the excellent oxidation resistance and content of ␣-SiC. On the one hand, oxidation at elevated temperatures may be contributed to decrease of flexural strength of SiBCN ceramics. SiBCN ceramics have excellent oxidation
resistance [3], though they contain BCN phase with a low resistance toward oxidative attack. The mass change of CS1900 and CD1900 were analyzed by TGA up to 1500 ◦ C in air, which shows no obvious mass change for CS1900 and 0.8% mass loss happened at 920 ◦ C for CD1900 (as shown in Fig. 5). These phenomena can also be found from the surface change of sam-
Fig. 5. TG spectra of CS1900 and CD1900 in air.
Z.-H. Yang et al. / Materials Science and Engineering A 489 (2008) 187–192
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Fig. 7. Thermal expansion coefficients of SiBCN ceramics.
contains higher content of 6H-SiC, is approximately higher at high temperature and increase more rapidly than that of CD1900. 4. Summary
Fig. 6. SEM surface images of the SiBCN ceramics after elevated temperature bending strength testing at 1000 ◦ C. (a) CS1900 and (b) CD1900.
ples tested at 1000 ◦ C for bending strength (as shown in Fig. 6). No visible change happened for the CS1900 at 1000 ◦ C, but uneven surface appeared for the CD1900 because of the oxidation of BCN phase. On the other hand, the lower reduction rate of strengths for CS1900 at elevated temperatures may be contributed to the high content of ␣-SiC. ␣-SiC is the most stable form of SiC at high temperatures [23]. Liu et al. [24] has investigated the effect of content of ␣-SiC on the hightemperature strength. They found that ␣/ (30:70) SiC ceramic and pure -SiC ceramic have the same flexural strengths at 1000 ◦ C. When the temperature increase to 1200 ◦ C, a slight decrease of strength by ∼9% for the ␣/ SiC ceramic happened. However, for pure -SiC ceramic, there is ∼17% reduction in strength as compared to the strength at 1000 ◦ C. So the high-temperature strengths of CS1900 have a lower reduction rate. The thermal expansion coefficients of CS1900 and CD1900 ceramics were almost same from room temperature to 1400 ◦ C (as shown in Fig. 7). The average values were 4.55 × 10−6 /◦ C and 4.46 × 10−6 /◦ C for CS1900 and CD1900, respectively, which were almost equal to values of pure SiC ceramics [25]. 6H-SiC have different coefficient of thermal expansion at different crystal direction. The coefficient of thermal expansion in the 2 1¯ 1¯ 0 direction, α11 , is different from the coefficient in the 0 0 0 1 direction, α33 . On the basis of measurements from −120 to 1400 ◦ C, Taylor and Jones [26] have reported below 750 ◦ C, α11 < α33 ; above 750 ◦ C the α11 value exceeds α33 So, the thermal expansion coefficient of CS1900 ceramic, which
In summary, two SiB0.5 C1.5 N0.5 ceramics were consolidated by hot pressing using one- and two-step milled powders. SiB0.5 C1.5 N0.5 ceramics had different microstructures, but both of ceramics mainly contained nano-sized ␣-SiC (6H), -SiC (3C) and BCN phase. For the SiB0.5 C1.5 N0.5 ceramics prepared by the one-step milled powders contained high content of ␣SiC. For the SiB0.5 C1.5 N0.5 ceramics prepared by the two-step milled powders mainly contained -SiC. Compared ceramics sintered at 1900 ◦ C, the former has higher fracture toughness lower than that of the latter. The flexural strength of former is 312.8 MPa at room temperature, and the strengths reduce rates are 8.2% and 28.0% at 1000 and 1400 ◦ C, respectively. For the latter, the flexural strength is 423 MPa at room temperature, but the strength decreased 11.3% at 1000 ◦ C and 38.5% at 1400 ◦ C. The low strengths reduce rates for former ceramics should be contributed to the excellent oxidation resistance and high content of ␣-SiC. The average thermal expansion coefficients of SiBCN ceramics were about 4.5 × 10−6 /◦ C, which were approximate equal to that of pure SiC ceramics. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (NSFC, Grant number 90505011) and Program for New Century Excellent Talents in University (NCET, Grant number NCET-04-0327). References [1] R. Riedel, A. Kienzle, W. Dressler, L. Ruwisch, J. Bill, F. Aldinger, Nature 382 (1996) 796. [2] N.V.R. Kumar, S. Prinz, Y. Cai, A. Zimmermann, F. Aldinger, F. Berger, K. M¨uller, Acta Mater. 53 (2005) 4567. [3] N.S. Jacobson, E.J. Opila, K.N. Lee, Curr. Opin. Solid State Mater. Sci. 5 (2001) 301. [4] A. Zimmermann, A. Bauer, M. Christ, Y. Cai, F. Aldinger, Acta Mater. 50 (2002) 1187. [5] N.V. Ravi Kumar, R. Mager, Y. Cai, A. Zimmermann, F. Aldinger, Scripta Mater. 51 (2004) 65.
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[6] B. Baufeld, H. Gu, J. Bill, F. Wakai, F. Aldinger, J. Eur. Ceram. Soc. 19 (1999) 2797. [7] R. Riedel, L.M. Ruswisch, L. An, R. Raj, J. Am. Ceram. Soc. 81 (1998) 3341. [8] D.L. Zhang, Prog. Mater. Sci. 49 (2004) 537. [9] X.Y. Yang, Z.W. Huang, Y.K. Wu, H.Q. Ye, Mater. Sci. Eng. A 300 (2001) 278. [10] X. Xie, Z. Yang, R. Ren, L.L. Shaw, Mater. Sci. Eng. A 255 (1998) 39. [11] Z.G. Yang, L.L. Shaw, Nanostruct. Mater. 7 (1996) 873. [12] M. Ohyanagi, T. Yamamoto, H. Kitaura, Y. Kodera, T. Ishii, Z.A. Munir, Scripta Mater. 50 (2004) 111. [13] L.L. Shaw, Z.-G. Yang, R.-M. Ren, Mater. Sci. Eng. A 244 (1998) 113. [14] B. Yao, L. Liu, W.H. Su, J. Appl. Phys. 86 (1999) 2464. [15] Y.F. Zhang, Y.H. Tang, C.S. Lee, I. Bello, S.T. Lee, Diam. Relat. Mater. 8 (1999) 610. [16] R. Torres, I. Caretti, R. Gago, Z. Mart´ın, I. Jim´enez, Diam. Relat. Mater. 16 (2007) 1450.
[17] Z.H. Yang, D.C. Jia, Y. Zhou, J.X. Zhang, Mater. Sci. Eng. A 488 (2008) 241–246. [18] Z.H. Yang, Y. Zhou, D.C. Jia, Nanotechnology, submitted for publication. [19] X. Xiao, Y. Li, L. Song, X. Peng, X. Hu, Appl. Surf. Sci. 156 (2000) 155. [20] Y.H. Xiong, C.S. Xiong, S.Q. Wei, H.W. Yang, Y.T. Mai, W. Xu, S. Yang, G.H. Dai, S.J. Song, J. Xiong, Z.M. Ren, J. Zhang, H.L. Pi, Z.C. Xia, S.L. Yuan, Appl. Surf. Sci. 253 (2006) 2515. [21] B. Maquin, A. Derr´e, C. Labrug`ere, M. Trinquecoste, P. Chadeyron, P. Delha`es, Carbon 38 (2000) 145–156. [22] G.D. Zhan, R.J. Xie, M.M.Y.W. Kim, J. Am. Ceram. Soc. 84 (2001) 945. [23] Y.W. Kim, M. Mitomo, H. Emoto, J.G. Lee, J. Am. Ceram. Soc. 81 (12) (1998) 3136–3140. [24] D.M. Liu, Z.C. Jou, B.W. Lin, C.T. Fu, J. Mater. Sci. Lett. 14 (1995)1327. [25] Z. Li, R.C. Bradt, J. Am. Ceram. Soc. 70 (1987) 445. [26] A. Taylor, R.M. Jones, in: J.R. O’Connor, J. Smiltens (Eds.), SiC-High Temperature Semiconductor, Pergamon Press, London, 1960, pp. 147–154.
Microstructures and properties of SiB0.5C1.5N0.5 ceramics consolidated by mechanical alloying and hot pressing Zhi-Hua Yang ∗ , Yu Zhou, De-Chang Jia, Qing-Chang Meng Institute for Advanced Ceramics, Harbin Institute of Technology, Harbin 150080, PR China Received 27 August 2007; received in revised form 7 December 2007; accepted 7 December 2007
Abstract Nanostructured SiB0.5 C1.5 N0.5 ceramics were formed by two routes of mechanical alloying and consolidated by hot pressing. Nano-sized 6H-SiC, 3C-SiC and BCN were the main phases for two SiB0.5 C1.5 N0.5 ceramics. For the SiBCN ceramics fabricated from the powders milled by one-step milling, the content of 6H-SiC is higher than that of ceramic fabricated from the powders milled by two-step milling. The flexural strengths for the former and latter were 312.8 MPa and 423.4 MPa at room temperature, respectively. At high temperature, the strengths of the former reduced 8.2% and 28.0% at 1000 ◦ C and 1400 ◦ C, respectively. The reduction rates of strengths for the latter were 11.3% and 38.5%, respectively. The average thermal expansion coefficients of two SiBCN ceramics consolidated at 1900 ◦ C were about 4.5 × 10−6 /◦ C. © 2008 Elsevier B.V. All rights reserved. Keywords: SiB0.5 C1.5 N0.5 ceramics; Mechanical alloying; Hot pressing; Mechanical properties; Microstructure
1. Introduction SiBCN ceramics derived from polymer precursors recently received a significant attention because of their outstanding hightemperature stability [1] and mechanical properties [2]. SiBCN ceramics had low density and the lowest reported oxidation rates of any non-oxide material known to date [3], so those ceramics suggest themselves for applications under extreme conditions like in heat engines. Some works have been done to develop such materials from the stage of laboratory curiosities to engineering materials by attempting industrial-scale processing, and, hence, it is very important to know the mechanical and thermal properties of SiBCN ceramics. However, up to now, only hightemperature compression creep properties were reported [4–7]. It may be due to the method of polymer precursors limiting the fabrications of big size bulk ceramics. Powder processing by mechanical alloying (MA) has attracted wide practical interest as it offers a simple but powerful way to synthesize non-equilibrium phases and microstructures, from nanograin materials to extended solid solutions, amorphous phases, chemically disordered compounds, and nanocomposites [8]. SiC [9–12], Si3 N4 –SiC [13], and BCN ∗
Corresponding author. Tel.: +86 451 86418792; fax: +86 451 86414291. E-mail address: [email protected] (Z.-H. Yang).
0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.12.010
[14–16] powders have been fabricated by MA. Our previous work established that nanostructured composite SiB0.5 C1.5 N0.5 can be produced by MA [17]. In the present paper, SiBCN ceramics were consolidated by hot pressing (HP) using two types of milled SiBCN powders as starting materials. Microstructures of SiBCN ceramics were characterized. Mechanical properties (at room and high temperature) and thermal expansion coefficient were also tested. 2. Experimental procedure SiB0.5 C1.5 N0.5 powders (with a target composition of SiC–33.3 at.% BCN) were fabricated by one- and two-step MA using shake mill [17,18]. For one-step milling, c-Si, h-BN and graphite (C) powders were simultaneously milled for 20 h. As for two-step milling, both Si and partial C powders in mole ratio of 1:1 were milled for 15 h, and then the residual C and BN powders were added into the vessel for another 5 h milling. The powders prepared by one- and two-step milling were denoted as CS and CD, respectively. Two SiB0.5 C1.5 N0.5 powders mainly contained amorphous SiBCN, nano-crystalline (4–5 nm) 3CSiC and 6H-SiC. In addition, about 4.7% Si remained in the CS powders. Sintering was performed by HP at 1850/1900 ◦ C under a pressure of 40 MPa for 30 min in N2 (1.0 × 105 Pa). The furnace for
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Fig. 1. XRD patterns of the bulk samples produced by HP of SiBCN powders at various holding temperatures.
tent with the analysis results of the milled SiBCN powders [18]. In addition, the appearance of ZrO2 reflections was due to the contamination from the wear of ZrO2 balls and inner surface of vessel. The relative density is 91.5% for the CS1850, and it can reach 99.4% for CS1900. However, the relative density of CD1900 is only 94.3%. SEM photographs of polishing surface of SiBCN ceramics are shown in Fig. 2. As can be seen that the size of all grains are <1 m. Fine, equiaxed SiC grains (light phase) were homogeneously distributed in SiBCN ceramics. Many micro cracks were found the around the SiC grains, and some cracks even crossed the SiC grains, as shown by the white circle in the Fig. 2c. Micro cracks may be caused by the thermal shock during the module transferred from the sintering chamber to depositing chamber. The TEM images of CS and CD powders consolidated at different temperatures are shown in Fig. 3. For the CS1850, SiC
the sintering has two chambers. One chamber is for the sintering, and the other one is used to deposit module which was transferred from the sintering chamber. The temperature of depositing chamber is below 100 ◦ C during the sintering. The module was transferred into the depositing chamber when the sintering temperature decreased to 900 ◦ C in order to shorten the duration of sintering cycle. Hereafter, the sintered samples were marked as CS/CD#, # = 1850 or 1900, corresponding to consolidation temperatures of 1850 and 1900 ◦ C. The resulting sample ceramics were cut into bars of 36 mm × 3 mm × 4 mm (30 mm outer span) for measuring flexural strength at room and high temperature with a crosshead speed of 0.5 mm/min. The samples were heated up to 1000 and 1400 ◦ C at a rate of 10 ◦ C/min in air, and held at chosen temperature for 10 min before loading to ensure thermal equilibrium. The fracture toughness was determined using the single edge notched beam method with a crosshead speed of 0.05 mm/min. The testing bar dimensions used were 2 mm × 4 mm × 20 mm (16 mm outer span). The depth of the notches was 2.0 mm and the width about 0.2 mm. The structural characterization of ceramics was analyzed using X-ray diffraction (XRD) methods with Cu K␣ radiation. The investigation of microstructure was done in scanning electron microscope (SEM) and transmission electron microscope (TEM) operating at 120 kV. Thermogravimetric analysis (TGA) of the ceramics was performed in Al2 O3 crucible (STA-449C, Netzsch) by heating up to 1500 ◦ C at a heating rate of 10 ◦ C/min in air. Thermal expansion coefficient was performed by heating up to 1400 ◦ C at a heating rate of 10 ◦ C/min in argon. 3. Results and discussion Fig. 1 shows the XRD patterns of the bulk samples produced by HP of SiBCN powders at various holding temperatures. In all cases, clear diffraction peaks are observed suggesting that most grains have recrystallized during the sintering process. BCN, 3CSiC and 6H-SiC are the main phase for the CS and CD ceramics. 3C-SiC is the main SiC phase in CD1900, but 6H-SiC is predominant phase for CS ceramics. The coexistence of 3C/6H-SiC and their relative contents for the two SiBCN ceramics were consis-
Fig. 2. SEM images of polishing surface of SiBCN ceramics hot-pressed at different temperature. (a) CS1850, (b) CS1900, and (c) CD1900.
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Fig. 3. TEM images of CS and CD powders sintered at different temperatures. (a) CS1850, (b) CS1900, and (c) CD1900.
grain sizes range from 100 nm to 300 nm (as shown in Fig. 3a). BCN mainly distribute the grains boundaries of SiC crystal. The grain sizes of BCN phase are very small, and it is difficult to display the BCN nano-crystal by the TEM. Grain sizes of SiC crystal increased up to 500 nm when the sintering temperature increasing 50 ◦ C. And the Selected Area Electron Diffraction (SAD) patterns from the selected area b-1 and b-2 exhibit [0 0 1] and [0 1 1¯ 0] patterns indicative of the -3C and ␣-6H structures. Some of BCN crystal visibly formed, and grain sizes increased for the CS1900 (as shown in Fig. 3b). For the CD1900, the grain sizes of -SiC are similar with these of CS1900 (as shown in Fig. 3c). BCN phase do not have visible crystals, but the size of BCN area is bigger than that in CS1850 and CS1900, which may be caused by the shorter times of milling (5 h) for BN-C powders in CD than these (20 h) in CS. The surface chemistry of CS1900 platelet was characterized by XPS. The measurement results are shown in Fig. 4. The characteristic peaks at 99.8 eV and 284.3 eV were identified as Si2p and C1s signals of the SiC covalent bond, respectively [19]. The
peaks at 191.0 eV (B1s), 285.2 eV (C1s), and 400.8 eV (N1s) can reflect the existence of B–C–N bond [20]. The peaks at 190.6 eV (B1s) and 397.7 eV (N1s) may be assigned to the B–N bonds of h-BN [20]. There is a peak at 284.3 eV for C1s spectrum, which assigned to the C–C bonds of graphite and agreed well with result of starting graphite [21]. The peak at 101.4 eV for Si2p spectrum was assigned to the Si–N bonds [19], but the content of Si–N bonds was <4% according the calculating method of SiBCN powders [18]. Therefore, it is difficult to detect by the XRD for the Si–N bonds. Structural characterization by XPS recorded the coexistence of several bonds between Si, B, C and N atoms. The synthesized SiBCN ceramics appeared to possess a structure containing several bonds such as Si–C, Si–N, B–C–N, C–C with B–N hybridization, rather than a simple complex of these bonds. Mechanical properties of SiBCN ceramics at room and high temperature are shown in Table 1. It can be seen that the sintering temperature and content of ␣-SiC affected mechanical properties of SiBCN ceramics. Mechanical properties of CS
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Fig. 4. (a) Si2p, (b) B1s, (c) C1s and (d) N1s XPS spectra, with deconvolutions of peaks, recorded for CS1900 ceramics. Table 1 Mechanical properties of SiBCN ceramics Sample
CS1850 CS1900 CD1900
Flexural strength (MPa) RT
1000 ◦ C
1400 ◦ C
191.7 ± 20.7 312.8 ± 5.2 423.4 ± 43.0
– 287.2 ± 15.5 375.4 ± 10.7
– 225.3 ± 15.1 260.4 ± 15.7
Elastic modulus (GPa)
Fracture toughness (MPa m1/2 )
Vickers hardness (GPa)
107.9 ± 9.5 136.3 ± 17.8 134.6 ± 13.5
1.80 ± 0.07 3.31 ± 0.02 3.09 ± 0.05
2.44 ± 0.31 4.17 ± 0.50 3.89 ± 0.43
ceramics increased with the increase of sintering temperature. The increase of properties could be contributed to the increase of relative density from 91.5% to 99.4%. As for the CS1900 and CD1900, the content of ␣-SiC have a great influence on the mechanical properties. The microstructure of SiC obtained from an ␣-SiC powder is composed of fine equiaxed grains and is usually brittle. And the ceramics, which have higher content of ␣-SiC, have higher fracture toughness [22]. So it can be seen that the flexural strength of CD1900 is 110 MPa higher than that of CS1900 at room temperature (RT), despite the low density of CD1900. The fracture toughness of CS1900 is higher than that of CD1900. CD1900 exhibits higher strengths at RT and high temperature than CS1900. High-temperature strengths of CD1900 reduced 11.3% at 1000 ◦ C and 38.5% at 1400 ◦ C, respectively, but the values were only 8.2% and 28.0% for CS1900. The reasons for the lower reduction rates of high-temperature strength of CS1900 may be related to the excellent oxidation resistance and content of ␣-SiC. On the one hand, oxidation at elevated temperatures may be contributed to decrease of flexural strength of SiBCN ceramics. SiBCN ceramics have excellent oxidation
resistance [3], though they contain BCN phase with a low resistance toward oxidative attack. The mass change of CS1900 and CD1900 were analyzed by TGA up to 1500 ◦ C in air, which shows no obvious mass change for CS1900 and 0.8% mass loss happened at 920 ◦ C for CD1900 (as shown in Fig. 5). These phenomena can also be found from the surface change of sam-
Fig. 5. TG spectra of CS1900 and CD1900 in air.
Z.-H. Yang et al. / Materials Science and Engineering A 489 (2008) 187–192
191
Fig. 7. Thermal expansion coefficients of SiBCN ceramics.
contains higher content of 6H-SiC, is approximately higher at high temperature and increase more rapidly than that of CD1900. 4. Summary
Fig. 6. SEM surface images of the SiBCN ceramics after elevated temperature bending strength testing at 1000 ◦ C. (a) CS1900 and (b) CD1900.
ples tested at 1000 ◦ C for bending strength (as shown in Fig. 6). No visible change happened for the CS1900 at 1000 ◦ C, but uneven surface appeared for the CD1900 because of the oxidation of BCN phase. On the other hand, the lower reduction rate of strengths for CS1900 at elevated temperatures may be contributed to the high content of ␣-SiC. ␣-SiC is the most stable form of SiC at high temperatures [23]. Liu et al. [24] has investigated the effect of content of ␣-SiC on the hightemperature strength. They found that ␣/ (30:70) SiC ceramic and pure -SiC ceramic have the same flexural strengths at 1000 ◦ C. When the temperature increase to 1200 ◦ C, a slight decrease of strength by ∼9% for the ␣/ SiC ceramic happened. However, for pure -SiC ceramic, there is ∼17% reduction in strength as compared to the strength at 1000 ◦ C. So the high-temperature strengths of CS1900 have a lower reduction rate. The thermal expansion coefficients of CS1900 and CD1900 ceramics were almost same from room temperature to 1400 ◦ C (as shown in Fig. 7). The average values were 4.55 × 10−6 /◦ C and 4.46 × 10−6 /◦ C for CS1900 and CD1900, respectively, which were almost equal to values of pure SiC ceramics [25]. 6H-SiC have different coefficient of thermal expansion at different crystal direction. The coefficient of thermal expansion in the 2 1¯ 1¯ 0 direction, α11 , is different from the coefficient in the 0 0 0 1 direction, α33 . On the basis of measurements from −120 to 1400 ◦ C, Taylor and Jones [26] have reported below 750 ◦ C, α11 < α33 ; above 750 ◦ C the α11 value exceeds α33 So, the thermal expansion coefficient of CS1900 ceramic, which
In summary, two SiB0.5 C1.5 N0.5 ceramics were consolidated by hot pressing using one- and two-step milled powders. SiB0.5 C1.5 N0.5 ceramics had different microstructures, but both of ceramics mainly contained nano-sized ␣-SiC (6H), -SiC (3C) and BCN phase. For the SiB0.5 C1.5 N0.5 ceramics prepared by the one-step milled powders contained high content of ␣SiC. For the SiB0.5 C1.5 N0.5 ceramics prepared by the two-step milled powders mainly contained -SiC. Compared ceramics sintered at 1900 ◦ C, the former has higher fracture toughness lower than that of the latter. The flexural strength of former is 312.8 MPa at room temperature, and the strengths reduce rates are 8.2% and 28.0% at 1000 and 1400 ◦ C, respectively. For the latter, the flexural strength is 423 MPa at room temperature, but the strength decreased 11.3% at 1000 ◦ C and 38.5% at 1400 ◦ C. The low strengths reduce rates for former ceramics should be contributed to the excellent oxidation resistance and high content of ␣-SiC. The average thermal expansion coefficients of SiBCN ceramics were about 4.5 × 10−6 /◦ C, which were approximate equal to that of pure SiC ceramics. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (NSFC, Grant number 90505011) and Program for New Century Excellent Talents in University (NCET, Grant number NCET-04-0327). References [1] R. Riedel, A. Kienzle, W. Dressler, L. Ruwisch, J. Bill, F. Aldinger, Nature 382 (1996) 796. [2] N.V.R. Kumar, S. Prinz, Y. Cai, A. Zimmermann, F. Aldinger, F. Berger, K. M¨uller, Acta Mater. 53 (2005) 4567. [3] N.S. Jacobson, E.J. Opila, K.N. Lee, Curr. Opin. Solid State Mater. Sci. 5 (2001) 301. [4] A. Zimmermann, A. Bauer, M. Christ, Y. Cai, F. Aldinger, Acta Mater. 50 (2002) 1187. [5] N.V. Ravi Kumar, R. Mager, Y. Cai, A. Zimmermann, F. Aldinger, Scripta Mater. 51 (2004) 65.
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