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Densification of high-strength B4C–TiB2 composites fabricated by pulsed electric current sintering of TiC–B mixture
Scripta Materialia 135 (2017) 15–18
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Densification of high-strength B4C–TiB2 composites fabricated by pulsed electric current sintering of TiC–B mixture Zetan Liu a,b, Dewen Wang c,d, Jiamao Li b, Qing Huang d, Songlin Ran a,b,⁎ a
Key Laboratory of Metallurgical Emission Reduction & Resources Recycling, Ministry of Education, Anhui University of Technology, 243002 Ma'anshan, China Anhui Key Laboratory of Metal Materials and Processing, School of Materials Science and Engineering, Anhui University of Technology, 243002 Ma'anshan, China State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Science, 200050 Shanghai, China d Engineering Laboratory of Specialty Fibers and Nuclear Energy Materials (FiNE), Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, 315201 Ningbo, China b c
a r t i c l e
i n f o
Article history: Received 5 February 2017 Accepted 16 March 2017 Keywords: Boron carbide (B4C) Titanium diboride (TiB2) Reactive pulsed electric current sintering Mechanical properties
a b s t r a c t The densification and mechanical properties of B4C–TiB2 composites fabricated by reactive pulsed electric current sintering from a mixture of TiC and amorphous B powders were investigated. The excess of B was essential to remove the carbon produced by the reaction. The degassing process at 1900 °C before applying a 50 MPa external pressure greatly improved the densification of the composites. The B4C–41 vol% TiB2 composite obtained at the optimum condition had a high 3-point bending strength of 891 MPa, a Vickers hardness of 28 GPa and a fracture toughness of 4.4 MPa m1/2, respectively. © 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
As one of the hardest and the lightest ceramics, B4C has becoming increasingly popular for a variety of structural applications. However, due to its covalent bond, the densification of B4C ceramic requires different additives, which normally weights and softens the material [1–4]. In addition, it is very difficult to machine the hard B4C ceramic into complex shapes using traditional diamond tools. Although B4C is semiconducting, its electrical resistance is still too high to satisfy the demand for the electrical discharge machining (EDM) [5]. B4C–TiB2 composite has been recently becoming one of the most potential B4C based composites, since the addition of TiB2 to B4C not only increases the densification but also decreases the electrical resistance as well as retains the extreme hardness and the low density of monolithic B4C [1,2]. Generally, B4C–TiB2 composites are prepared by pressureless sintering (PLS), hot-pressed sintering (HPS) or pulsed electric current sintering (PECS) from commercially available B4C and TiB2 powders [3,6–8]. Among these sintering techniques, PECS, also known as spark plasma sintering (SPS) or plasma activation sintering (PAS) exhibits remarkable advantages over densification and grain growth inhibition. Fully dense B4C–40 vol% TiB2 composites prepared by PECS were reported with a B4C grain size of 1.57 μm combining an excellent hardness of 31 GPa and a flexural strength above 800 MPa [3]. ⁎ Corresponding author at: Key Laboratory of Metallurgical Emission Reduction & Resources Recycling, Ministry of Education, Anhui University of Technology, 243002 Ma'anshan, China. E-mail address: [email protected] (S. Ran).
http://dx.doi.org/10.1016/j.scriptamat.2017.03.023 1359-6462/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
As an alternative route, B4C–TiB2 composites prepared via in situ reactive PS, HPS or PECS were reported to have fine microstructures and excellent mechanical properties [4,9–15]. For example, B4C–30 vol% TiB2 composite prepared by reactive PECS has a high hardness of 39.3 GPa and a flexural strength of 865 MPa [4], and B4C–43 vol% TiB2 composite prepared by co-precipitation and reactive HPS has a fracture toughness as high as 9.4 MPa m1/2 [10], respectively. Among these report [4,9–15], B4C–TiB2 composites were synthesized from either elemental powders or the reaction among B4C, TiO2 and carbon. Recently, we successfully prepared B4C–TiB2 composite powders via a carbide boronizing process according to reaction (1) on the basis of an early report [1,2,16]. The purpose of this paper is to report on the fabrication and properties of bulk B4C–TiB2 composites via reactive PECS with the same carbide boronizing process. TiC þ 6B ¼ B4 C þ TiB2
ð1Þ
Commercially available TiC (purity N99.0%, 3.0 μm) and amorphous boron powders (purity 95.8%, 0.90 μm) were ball mixed in ethanol for 24 h using ZrO2 beads. The molar ratio of TiC/B was set as 1:6 and 1:6.6. After mixing, the slurry was dried at 65 °C in a rotating evaporator and subsequently sieved with a 200 mesh sieve to minimize powder segregation and agglomeration. The synthesis process was conducted in a pulsed electric current sintering (PECS) furnace (Type HP D25, FCT System, Rauenstein, Germany) with a dynamic vacuum. In order to ensure constant contact of the electrodes with the die/punch set-up, a force of 5 kN was always applied
16
Z. Liu et al. / Scripta Materialia 135 (2017) 15–18
during the heating and dwelling process. The heating rate was 100 °C/min and a pressure of 50 MPa were applied in 60 s after the temperature reached 1900 °C for 0–3 min. After a dwell time of 4–10 min, the pressure was removed in 5 min while the temperature was rapidly cooled around 800 °C. When the temperature was below 400 °C, the samples were taken out after another 5 min. The obtained bulk material was about 3.0 mm in thickness and 40.0 mm in diameter. The disk was ground to remove the surface layer with a final thickness of 1.5 mm. The bulk density was measured by the Archimedes method. The phase composition was characterized by Xray diffraction (XRD, Bruker AXS D8 Discover, Germany). The microstructure of the ceramic was examined by a field scanning electron microscopy (FESEM, S-4800, Hitachi, Japan). The Vickers hardness was measured on a hardness tester with a load of 5 kg and a dwell time of 10 s. The indentation toughness was evaluated according to the indentation crack lengths. The reported values were the mean and standard deviations from 10 indentations. The flexural strength was measured by a 3-point bending test using specimens with dimensions 2.0 mm × 1.5 mm × 25 mm on a universal testing machine (Exceed E44, MTS China) with a span of 20 mm and a loading speed of 0.1 mm/min. The reported values were the mean and standard deviations of the best 5 of at least 10 bending bars. Fig. 1 presents XRD patterns of the bulk ceramics. For both patterns, B4C and TiB2 are the main crystalline phases, indicating B4C–TiB2 ceramic composites were successfully fabricated by reactive PECS from TiC and B powders. Although the theoretical volume percentage of B4C in B4C–TiB2 composite fabricated according to reaction (1) is as high as 59%, due to its lower atom number, its XRD peaks are much weaker than those of TiB2 phase. When the molar ratio of TiC/B is 1:6, a XRD peak for carbon is detected at 2Ɵ = 26.7°, as shown in Fig. 1(a), which is attributed to the proceeding of reaction (2). It is inevitable for commercial TiC and amorphous B powders to be oxidized on their surfaces in air since both of them are non-oxides. As a result, B was deficient in the reactive sintering process if the molar ratio of TiC/B was stoichiometric (1:6) according to reaction (1), and reaction (2) proceeded in the sintering process, which needs less B than reaction (1). The presence of carbon in ceramics is harmful for the hardness and strength of the material, and therefore it should be eliminated. When the molar ratio of TiC/B decreases to 1:6.6, the XRD peak for carbon disappeares, as shown in Fig. 1(b), revealing that the excess of B is effective in removal of carbon. In addition, the excess of B has beneficial effects on densification and hardness improvement of B4C based composites due to the formation of BxC (x N 4) [17]. Therefore, in this study, B4C–TiB2 composites
Fig. 1. XRD pattern of B4C–TiB2 ceramic fabricated by reactive PECS with a TiC/B molar ratio of (a) 1:6 and (b) 1:6.6.
were prepared with TiC/B = 1:6.6. TiC þ 2B ¼ TiB2 þ C
ð2Þ
Due to spontaneous oxidation of TiC and B, there was small amount of TiO2 and B2O3 in the system. During the heating process, TiO2 could be reduced to form TiB2 according to reactions (3) and produced B2O3. It is known that B2O3 has adverse effects on the densification of B4C or TiB2 [3,18,19]. B2O3 has a high vapor pressure above 1000 °C, it could completely evaporate in vacuum around 1450 °C during PLS process [20]. However, in the case of PECS, the result is different. The heating rate of PECS is much fast than that of PLS, the completion of B2O3 evaporation will be delayed to a higher temperature [3,21]. Moreover, as described in the part of experimental process, a force of 5 kN was always applied to keep the contact between electrodes and the die/punch setup, which isolated the reactive system from vacuum environment. Our previous research on fabrication of ZrB2-SiC composites by reactive PECS indicated that the oxide impurities could be completely removed by evaporation if there was a degassing time at sintering temperature before the final maximum pressure was applied [21]. Soon afterwards Huang et al. reported the effects of pressure-loading cycles on the densification of B4C–TiB2 composites by PECS and confirmed our conclusion [3,4]. 3TiO2 þ 10B ¼ 3TiB2 þ 2B2 O3 ðlÞ
ð3Þ
In this study, the effect of degassing time was investigated. Fig. 2 shows the sintering behaviour of TiC–B mixture with different degassing time at 1900 °C. Three parts with temperature fluctuations are found on the temperature curve during the heating process, as indicated by the circles in Fig. 2(a). The first part between 400 °C and 600 °C is due to the transition of an uncontrolled heating process to a controlled thermal cycling since the temperature was monitored and controlled by an optical pyrometer only when the temperature was above 400 °C. The third part around 1900 °C is attributed to the transition of a heating process to a dwelling process. The applying of pressure also contributed to the third temperature fluctuation in Fig. 2(a), which was delayed for 1–3 min in Fig. 2(b, c) since there was corresponding time for degassing. The second part begins at 1227 °C with a sharp increasing shrinkage, which was ascribed to exothermic reaction between TiC and B. In another our previous research, B4C–TiB2 composite powders were synthesized by pulsed electric current heating a mixture of TiC nanopowders and amorphous B powders using a special designed graphite die/punch set-up [1]. The reaction between TiC and B completed at 1000 °C without any dwelling time. It should be pointed out that the nanopowders have much higher activity than micropowders used in this study [2]. Therefore, the reaction temperature was much higher than that in reference [1]. The gas release is found to initiate at 1657 °C and reach a maximum value at 1827 °C, as shown in Fig. 1(a). If there is no extra degassing time (Fig. 1(a)), the value of the vacuum pressure is still in the range of the peak when the external pressure was applied. As a result, the gases were entrapped in the powder compact, which would subsequently deteriorate the densification of the ceramic in dwelling process. When there is 1 min for degassing (Fig. 2(b)), the value of the vacuum pressure decreases sharply to the border of the peak. With further increase of degassing time, as shown in Fig. 2(c), the value of the vacuum pressure is out of the peak. It is found that the applying of external pressure induced a further decrease of the vacuum pressure in Fig. 2(a, b) but not in Fig. 2(c), indicating the completion of gas evaporation after the degassing process. The effect of degassing time on the densification of ceramics was evaluated by their measured densities, as shown in Table 1. When there was no degassing process, the relative density (R. D.) of the ceramic prepared at 1900 °C is only 96.1%. After 1–3 min degassing, the densification of the ceramic was improved and reached a maximum R. D. value of 97.9% at 1 min. As described above, the gas evaporation was nearly completed after 1 min degassing, the more degassing time
Z. Liu et al. / Scripta Materialia 135 (2017) 15–18
17
Fig. 3. Back scatting electron images of a B4C–TiB2 composite (BT1900-1-7): (a) polished surface and (b) fractured surface.
Fig. 2. Sintering behaviour of TiC–B powder mixture (TiC/B = 1:6.6) at 1900 °C for 7 min under a pressure of 50 MPa with a degassing time of (a) 0, (b) 1, (c) 3 and (d) 5 min.
had no effect on the gas release but improved the grains growth since the powder compact was at a high temperature of 1900 °C without the external pressure. Table 1 shows the mechanical properties of as-prepared B4C–TiB2 composites. High densities induced high hardness and high strengthes. The Vickers' hardness of the as-prepared B4C–TiB2 composites is around 28 GPa. The maximum 3-point bending strength of B4C–41 vol% TiB2 composites in this study is 891 MPa. As a comparison, the highest reported strengths for B4C–TiB2 composites was 865 MPa (3-point bending) with 30 vol% TiB2 and 866 MPa (4-point bending) with 20 mol%
TiB2, respectively [4,12]. It should be noticed that all of the above B4C– TiB2 composites with high strength were prepared by either reactive PECS or reactive HPS, suggesting that the reactive sintering has great advantages in fabricating ceramic composites. The fracture toughness values are 4.4–5.0 MPa m1/2, which is in excellent agreement with that (4.4 MPa m1/2) reported in a B4C–40 vol% TiB2 composite prepared by PECS [3]. Ceramics with lower densities have more pores in the microstructures and thus result in higher toughness values. B4C–TiB2 composites are in situ synthesized during the sintering process with the reaction between TiC and B. Unlike those reported in the reference where only TiB2 was in situ produced [4,12], B4C and TiB2 are simultaneously generated in this study. Due to the pinning effect of freshly produced TiB2 and B4C particles, the ceramic has a more homogenous and finer microstructure, as shown in Fig. 3. The average grain sizes of both B4C (grey) and TiB2 (white) are b1 μm, much smaller than those reported in the literature [3]. In our recent research [1], due to the use of TiC nanopowders, the average grain size of both B4C and TiB2 in a B4C–TiB2 composite was reported to be only 0.5 μm. However, the relative density of the ceramic was only 96% even with a higher pressure (60 MPa). Therefore, the densification of B4C–TiB2 composites from nanopowders needs further investigations, which has great potential in mechanical improvement with ultrafine microstructures. In summary, B4C–TiB2 composites were in situ synthesized and densified at 1900 °C under a pressure of 50 MPa via a reactive pulsed electric current sintering technique from a mixture of TiC and amorphous B powders. Due to the spontaneous oxidation of raw powders, boron was deficient for the reaction and resulted in a small amount of carbon
Table 1 PECS parameters and properties of as-prepared B4C–TiB2 composites. Designation
Degassing time (min)
Dwell time (min)
Relative density (%)
Vickers hardness (GPa)
Bending strength (MPa)
Fracture toughness (MPa m1/2)
BT1900-0-7 BT1900-1-7 BT1900-3-7
0 1 3
7 7 7
96.1 97.9 96.9
27.2 ± 1.2 28.3 ± 1.4 28.3 ± 1.0
778 ± 12 891 ± 32 843 ± 32
5.0 ± 0.8 4.4 ± 0.5 5.0 ± 0.2
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impurity in the composite, which was solved by the excess of B. The effects of degassing time on the densification and mechanical properties of obtained B4C–TiB2 composites were investigated. The optimum degassing time was 1 min and the best sample exhibited a relative density of 97.9%, a 3-point bending strength of 891 MPa, a Vickers hardness of 28 GPa and a fracture toughness of 4.4 MPa m1/2, respectively. The high strength was attributed to the fine microstructure and homogeneous distribution induced by the pinning effects of in situ synthesized B4C and TiB2 grains. This work was supported by the National Natural Science Foundation of China (No. 51302002, 51502309). Zetan Liu thanks the support from the Graduate Innovation Foundation of Anhui University of Technology (No. 2015075). References [1] D. Wang, S. Ran, L. Shen, H. Sun, Q. Huang, J. Eur. Ceram. Soc. 35 (3) (2015) 1107–1112. [2] D. Wang, H. Sun, Q. Deng, Z. Ding, S. Ran, Q. Huang, Ceram. Int. 40 (9) (2014) 15341–15344. [3] S.G. Huang, K. Vanmeensel, O.J.A. Malek, O. Van der Biest, J. Vleugels, Mater. Sci. Eng. A 528 (3) (2011) 1302–1309.
[4] S.G. Huang, K. Vanmeensel, O. Van der Biest, J. Vleugels, J. Eur. Ceram. Soc. 31 (4) (2011) 637–644. [5] O. Malek, J. Vleugels, K. Vanmeensel, S. Huang, J. Liu, S. Van den Berghe, A. Datye, K.-H. Wu, B. Lauwers, J. Eur. Ceram. Soc. 31 (11) (2011) 2023–2030. [6] H. Baharvandi, A. Hadian, J. Mater. Eng. Perform. 17 (6) (2008) 838–841. [7] H.R. Baharvandi, A. Hadian, A. Alizadeh, Appl. Compos. Mater. 13 (3) (2006) 191–198. [8] L.S. Sigl, H.-J. Kleebe, J. Am. Ceram. Soc. 78 (9) (1995) 2374–2380. [9] D. Dudina, D. Hulbert, D. Jiang, C. Unuvar, S. Cytron, A. Mukherjee, J. Mater. Sci. 43 (10) (2008) 3569–3576. [10] X. Yue, S. Zhao, P. Lü, Q. Chang, H. Ru, Mater. Sci. Eng. A 527 (27–28) (2010) 7215–7219. [11] V. Skorokhod, M. Vlajic, V. Krstic, J. Mater. Sci. Lett. 15 (15) (1996) 1337–1339. [12] S. Yamada, K. Hirao, Y. Yamauchi, S. Kanzaki, J. Eur. Ceram. Soc. 23 (7) (2003) 1123–1130. [13] V. Skorokhod, V.D. Krstic, J. Mater. Sci. Lett. 19 (3) (2000) 237–239. [14] V.V. Skorokhod, Powder Metall. Met. Ceram. 39 (9) (2000) 504–513. [15] V.V. Skorokhod, Powder Metall. Met. Ceram. 39 (7) (2000) 414–423. [16] H. Itoh, K. Sugiura, H. Iwahara, J. Alloys Compd. 232 (1–2) (1996) 186–191. [17] J. Zou, S.G. Huang, K. Vanmeensel, G.J. Zhang, J. Vleugels, O. Biest, J. Am. Ceram. Soc. 96 (4) (2013) 1055–1059. [18] S.L. Dole, S. Prochazka, R.H. Doremus, J. Am. Ceram. Soc. 72 (6) (1989) 958–966. [19] S. Baik, P.F. Becher, J. Am. Ceram. Soc. 70 (8) (1987) 527–530. [20] S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, J. Am. Ceram. Soc. 89 (5) (2006) 1544–1550. [21] S. Ran, O. Van der Biest, J. Vleugels, J. Eur. Ceram. Soc. 30 (12) (2010) 2633–2642.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Densification of high-strength B4C–TiB2 composites fabricated by pulsed electric current sintering of TiC–B mixture Zetan Liu a,b, Dewen Wang c,d, Jiamao Li b, Qing Huang d, Songlin Ran a,b,⁎ a
Key Laboratory of Metallurgical Emission Reduction & Resources Recycling, Ministry of Education, Anhui University of Technology, 243002 Ma'anshan, China Anhui Key Laboratory of Metal Materials and Processing, School of Materials Science and Engineering, Anhui University of Technology, 243002 Ma'anshan, China State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Science, 200050 Shanghai, China d Engineering Laboratory of Specialty Fibers and Nuclear Energy Materials (FiNE), Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, 315201 Ningbo, China b c
a r t i c l e
i n f o
Article history: Received 5 February 2017 Accepted 16 March 2017 Keywords: Boron carbide (B4C) Titanium diboride (TiB2) Reactive pulsed electric current sintering Mechanical properties
a b s t r a c t The densification and mechanical properties of B4C–TiB2 composites fabricated by reactive pulsed electric current sintering from a mixture of TiC and amorphous B powders were investigated. The excess of B was essential to remove the carbon produced by the reaction. The degassing process at 1900 °C before applying a 50 MPa external pressure greatly improved the densification of the composites. The B4C–41 vol% TiB2 composite obtained at the optimum condition had a high 3-point bending strength of 891 MPa, a Vickers hardness of 28 GPa and a fracture toughness of 4.4 MPa m1/2, respectively. © 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
As one of the hardest and the lightest ceramics, B4C has becoming increasingly popular for a variety of structural applications. However, due to its covalent bond, the densification of B4C ceramic requires different additives, which normally weights and softens the material [1–4]. In addition, it is very difficult to machine the hard B4C ceramic into complex shapes using traditional diamond tools. Although B4C is semiconducting, its electrical resistance is still too high to satisfy the demand for the electrical discharge machining (EDM) [5]. B4C–TiB2 composite has been recently becoming one of the most potential B4C based composites, since the addition of TiB2 to B4C not only increases the densification but also decreases the electrical resistance as well as retains the extreme hardness and the low density of monolithic B4C [1,2]. Generally, B4C–TiB2 composites are prepared by pressureless sintering (PLS), hot-pressed sintering (HPS) or pulsed electric current sintering (PECS) from commercially available B4C and TiB2 powders [3,6–8]. Among these sintering techniques, PECS, also known as spark plasma sintering (SPS) or plasma activation sintering (PAS) exhibits remarkable advantages over densification and grain growth inhibition. Fully dense B4C–40 vol% TiB2 composites prepared by PECS were reported with a B4C grain size of 1.57 μm combining an excellent hardness of 31 GPa and a flexural strength above 800 MPa [3]. ⁎ Corresponding author at: Key Laboratory of Metallurgical Emission Reduction & Resources Recycling, Ministry of Education, Anhui University of Technology, 243002 Ma'anshan, China. E-mail address: [email protected] (S. Ran).
http://dx.doi.org/10.1016/j.scriptamat.2017.03.023 1359-6462/© 2017 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
As an alternative route, B4C–TiB2 composites prepared via in situ reactive PS, HPS or PECS were reported to have fine microstructures and excellent mechanical properties [4,9–15]. For example, B4C–30 vol% TiB2 composite prepared by reactive PECS has a high hardness of 39.3 GPa and a flexural strength of 865 MPa [4], and B4C–43 vol% TiB2 composite prepared by co-precipitation and reactive HPS has a fracture toughness as high as 9.4 MPa m1/2 [10], respectively. Among these report [4,9–15], B4C–TiB2 composites were synthesized from either elemental powders or the reaction among B4C, TiO2 and carbon. Recently, we successfully prepared B4C–TiB2 composite powders via a carbide boronizing process according to reaction (1) on the basis of an early report [1,2,16]. The purpose of this paper is to report on the fabrication and properties of bulk B4C–TiB2 composites via reactive PECS with the same carbide boronizing process. TiC þ 6B ¼ B4 C þ TiB2
ð1Þ
Commercially available TiC (purity N99.0%, 3.0 μm) and amorphous boron powders (purity 95.8%, 0.90 μm) were ball mixed in ethanol for 24 h using ZrO2 beads. The molar ratio of TiC/B was set as 1:6 and 1:6.6. After mixing, the slurry was dried at 65 °C in a rotating evaporator and subsequently sieved with a 200 mesh sieve to minimize powder segregation and agglomeration. The synthesis process was conducted in a pulsed electric current sintering (PECS) furnace (Type HP D25, FCT System, Rauenstein, Germany) with a dynamic vacuum. In order to ensure constant contact of the electrodes with the die/punch set-up, a force of 5 kN was always applied
16
Z. Liu et al. / Scripta Materialia 135 (2017) 15–18
during the heating and dwelling process. The heating rate was 100 °C/min and a pressure of 50 MPa were applied in 60 s after the temperature reached 1900 °C for 0–3 min. After a dwell time of 4–10 min, the pressure was removed in 5 min while the temperature was rapidly cooled around 800 °C. When the temperature was below 400 °C, the samples were taken out after another 5 min. The obtained bulk material was about 3.0 mm in thickness and 40.0 mm in diameter. The disk was ground to remove the surface layer with a final thickness of 1.5 mm. The bulk density was measured by the Archimedes method. The phase composition was characterized by Xray diffraction (XRD, Bruker AXS D8 Discover, Germany). The microstructure of the ceramic was examined by a field scanning electron microscopy (FESEM, S-4800, Hitachi, Japan). The Vickers hardness was measured on a hardness tester with a load of 5 kg and a dwell time of 10 s. The indentation toughness was evaluated according to the indentation crack lengths. The reported values were the mean and standard deviations from 10 indentations. The flexural strength was measured by a 3-point bending test using specimens with dimensions 2.0 mm × 1.5 mm × 25 mm on a universal testing machine (Exceed E44, MTS China) with a span of 20 mm and a loading speed of 0.1 mm/min. The reported values were the mean and standard deviations of the best 5 of at least 10 bending bars. Fig. 1 presents XRD patterns of the bulk ceramics. For both patterns, B4C and TiB2 are the main crystalline phases, indicating B4C–TiB2 ceramic composites were successfully fabricated by reactive PECS from TiC and B powders. Although the theoretical volume percentage of B4C in B4C–TiB2 composite fabricated according to reaction (1) is as high as 59%, due to its lower atom number, its XRD peaks are much weaker than those of TiB2 phase. When the molar ratio of TiC/B is 1:6, a XRD peak for carbon is detected at 2Ɵ = 26.7°, as shown in Fig. 1(a), which is attributed to the proceeding of reaction (2). It is inevitable for commercial TiC and amorphous B powders to be oxidized on their surfaces in air since both of them are non-oxides. As a result, B was deficient in the reactive sintering process if the molar ratio of TiC/B was stoichiometric (1:6) according to reaction (1), and reaction (2) proceeded in the sintering process, which needs less B than reaction (1). The presence of carbon in ceramics is harmful for the hardness and strength of the material, and therefore it should be eliminated. When the molar ratio of TiC/B decreases to 1:6.6, the XRD peak for carbon disappeares, as shown in Fig. 1(b), revealing that the excess of B is effective in removal of carbon. In addition, the excess of B has beneficial effects on densification and hardness improvement of B4C based composites due to the formation of BxC (x N 4) [17]. Therefore, in this study, B4C–TiB2 composites
Fig. 1. XRD pattern of B4C–TiB2 ceramic fabricated by reactive PECS with a TiC/B molar ratio of (a) 1:6 and (b) 1:6.6.
were prepared with TiC/B = 1:6.6. TiC þ 2B ¼ TiB2 þ C
ð2Þ
Due to spontaneous oxidation of TiC and B, there was small amount of TiO2 and B2O3 in the system. During the heating process, TiO2 could be reduced to form TiB2 according to reactions (3) and produced B2O3. It is known that B2O3 has adverse effects on the densification of B4C or TiB2 [3,18,19]. B2O3 has a high vapor pressure above 1000 °C, it could completely evaporate in vacuum around 1450 °C during PLS process [20]. However, in the case of PECS, the result is different. The heating rate of PECS is much fast than that of PLS, the completion of B2O3 evaporation will be delayed to a higher temperature [3,21]. Moreover, as described in the part of experimental process, a force of 5 kN was always applied to keep the contact between electrodes and the die/punch setup, which isolated the reactive system from vacuum environment. Our previous research on fabrication of ZrB2-SiC composites by reactive PECS indicated that the oxide impurities could be completely removed by evaporation if there was a degassing time at sintering temperature before the final maximum pressure was applied [21]. Soon afterwards Huang et al. reported the effects of pressure-loading cycles on the densification of B4C–TiB2 composites by PECS and confirmed our conclusion [3,4]. 3TiO2 þ 10B ¼ 3TiB2 þ 2B2 O3 ðlÞ
ð3Þ
In this study, the effect of degassing time was investigated. Fig. 2 shows the sintering behaviour of TiC–B mixture with different degassing time at 1900 °C. Three parts with temperature fluctuations are found on the temperature curve during the heating process, as indicated by the circles in Fig. 2(a). The first part between 400 °C and 600 °C is due to the transition of an uncontrolled heating process to a controlled thermal cycling since the temperature was monitored and controlled by an optical pyrometer only when the temperature was above 400 °C. The third part around 1900 °C is attributed to the transition of a heating process to a dwelling process. The applying of pressure also contributed to the third temperature fluctuation in Fig. 2(a), which was delayed for 1–3 min in Fig. 2(b, c) since there was corresponding time for degassing. The second part begins at 1227 °C with a sharp increasing shrinkage, which was ascribed to exothermic reaction between TiC and B. In another our previous research, B4C–TiB2 composite powders were synthesized by pulsed electric current heating a mixture of TiC nanopowders and amorphous B powders using a special designed graphite die/punch set-up [1]. The reaction between TiC and B completed at 1000 °C without any dwelling time. It should be pointed out that the nanopowders have much higher activity than micropowders used in this study [2]. Therefore, the reaction temperature was much higher than that in reference [1]. The gas release is found to initiate at 1657 °C and reach a maximum value at 1827 °C, as shown in Fig. 1(a). If there is no extra degassing time (Fig. 1(a)), the value of the vacuum pressure is still in the range of the peak when the external pressure was applied. As a result, the gases were entrapped in the powder compact, which would subsequently deteriorate the densification of the ceramic in dwelling process. When there is 1 min for degassing (Fig. 2(b)), the value of the vacuum pressure decreases sharply to the border of the peak. With further increase of degassing time, as shown in Fig. 2(c), the value of the vacuum pressure is out of the peak. It is found that the applying of external pressure induced a further decrease of the vacuum pressure in Fig. 2(a, b) but not in Fig. 2(c), indicating the completion of gas evaporation after the degassing process. The effect of degassing time on the densification of ceramics was evaluated by their measured densities, as shown in Table 1. When there was no degassing process, the relative density (R. D.) of the ceramic prepared at 1900 °C is only 96.1%. After 1–3 min degassing, the densification of the ceramic was improved and reached a maximum R. D. value of 97.9% at 1 min. As described above, the gas evaporation was nearly completed after 1 min degassing, the more degassing time
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Fig. 3. Back scatting electron images of a B4C–TiB2 composite (BT1900-1-7): (a) polished surface and (b) fractured surface.
Fig. 2. Sintering behaviour of TiC–B powder mixture (TiC/B = 1:6.6) at 1900 °C for 7 min under a pressure of 50 MPa with a degassing time of (a) 0, (b) 1, (c) 3 and (d) 5 min.
had no effect on the gas release but improved the grains growth since the powder compact was at a high temperature of 1900 °C without the external pressure. Table 1 shows the mechanical properties of as-prepared B4C–TiB2 composites. High densities induced high hardness and high strengthes. The Vickers' hardness of the as-prepared B4C–TiB2 composites is around 28 GPa. The maximum 3-point bending strength of B4C–41 vol% TiB2 composites in this study is 891 MPa. As a comparison, the highest reported strengths for B4C–TiB2 composites was 865 MPa (3-point bending) with 30 vol% TiB2 and 866 MPa (4-point bending) with 20 mol%
TiB2, respectively [4,12]. It should be noticed that all of the above B4C– TiB2 composites with high strength were prepared by either reactive PECS or reactive HPS, suggesting that the reactive sintering has great advantages in fabricating ceramic composites. The fracture toughness values are 4.4–5.0 MPa m1/2, which is in excellent agreement with that (4.4 MPa m1/2) reported in a B4C–40 vol% TiB2 composite prepared by PECS [3]. Ceramics with lower densities have more pores in the microstructures and thus result in higher toughness values. B4C–TiB2 composites are in situ synthesized during the sintering process with the reaction between TiC and B. Unlike those reported in the reference where only TiB2 was in situ produced [4,12], B4C and TiB2 are simultaneously generated in this study. Due to the pinning effect of freshly produced TiB2 and B4C particles, the ceramic has a more homogenous and finer microstructure, as shown in Fig. 3. The average grain sizes of both B4C (grey) and TiB2 (white) are b1 μm, much smaller than those reported in the literature [3]. In our recent research [1], due to the use of TiC nanopowders, the average grain size of both B4C and TiB2 in a B4C–TiB2 composite was reported to be only 0.5 μm. However, the relative density of the ceramic was only 96% even with a higher pressure (60 MPa). Therefore, the densification of B4C–TiB2 composites from nanopowders needs further investigations, which has great potential in mechanical improvement with ultrafine microstructures. In summary, B4C–TiB2 composites were in situ synthesized and densified at 1900 °C under a pressure of 50 MPa via a reactive pulsed electric current sintering technique from a mixture of TiC and amorphous B powders. Due to the spontaneous oxidation of raw powders, boron was deficient for the reaction and resulted in a small amount of carbon
Table 1 PECS parameters and properties of as-prepared B4C–TiB2 composites. Designation
Degassing time (min)
Dwell time (min)
Relative density (%)
Vickers hardness (GPa)
Bending strength (MPa)
Fracture toughness (MPa m1/2)
BT1900-0-7 BT1900-1-7 BT1900-3-7
0 1 3
7 7 7
96.1 97.9 96.9
27.2 ± 1.2 28.3 ± 1.4 28.3 ± 1.0
778 ± 12 891 ± 32 843 ± 32
5.0 ± 0.8 4.4 ± 0.5 5.0 ± 0.2
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impurity in the composite, which was solved by the excess of B. The effects of degassing time on the densification and mechanical properties of obtained B4C–TiB2 composites were investigated. The optimum degassing time was 1 min and the best sample exhibited a relative density of 97.9%, a 3-point bending strength of 891 MPa, a Vickers hardness of 28 GPa and a fracture toughness of 4.4 MPa m1/2, respectively. The high strength was attributed to the fine microstructure and homogeneous distribution induced by the pinning effects of in situ synthesized B4C and TiB2 grains. This work was supported by the National Natural Science Foundation of China (No. 51302002, 51502309). Zetan Liu thanks the support from the Graduate Innovation Foundation of Anhui University of Technology (No. 2015075). References [1] D. Wang, S. Ran, L. Shen, H. Sun, Q. Huang, J. Eur. Ceram. Soc. 35 (3) (2015) 1107–1112. [2] D. Wang, H. Sun, Q. Deng, Z. Ding, S. Ran, Q. Huang, Ceram. Int. 40 (9) (2014) 15341–15344. [3] S.G. Huang, K. Vanmeensel, O.J.A. Malek, O. Van der Biest, J. Vleugels, Mater. Sci. Eng. A 528 (3) (2011) 1302–1309.
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