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HfB2–SiC composite prepared by reactive spark plasma sintering
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CERAMICS INTERNATIONAL
Ceramics International 40 (2014) 11009–11013 www.elsevier.com/locate/ceramint
HfB2–SiC composite prepared by reactive spark plasma sintering Hailong Wanga, Sea-Hoon Leeb,n, Lun Fengb a School of Materials Science and Engineering, Zhengzhou University, Zhengzhou 450001, China Division of Powder/Ceramics Research, Korea Institute of Materials Science, 531 Changwondaero, Changwon, Gyeongnam 641–831, Republic of Korea
b
Received 10 March 2014; received in revised form 19 March 2014; accepted 19 March 2014 Available online 27 March 2014
Abstract Dense HfB2–SiC composites were fabricated by reactive spark plasma sintering of a powder mixture composed of HfSi2, B4C and carbon. The relative density of the composites was 98.7% after reactive spark plasma sintering at 1600 1C under a pressure of 40 MPa for 10 min. Strong exothermic reaction occurred during the heating of the powder mixture due to the following reaction: 2HfSi2 þB4Cþ3C ¼2HfB2 þ 4SiC. The in situ formed HfB2 and SiC phases were homogeneously distributed. The average grain size of formed HfB2 and SiC was about 2 and 1 μm, respectively. The Vickers hardness and fracture toughness of HfB2–SiC composites were 20.47 1.5 GPa and 4.7 70.3 MPa m1/2, respectively. The present report proposes a method to densify HfB2-based composites at a low temperature. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: Hafnium diboride; Microstructure; Reactive spark plasma sintering
1. Introduction Hafnium diboride (HfB2) is a potential candidate for ultra-hightemperature applications in the aerospace industry because of its high melting point, high strength, high thermal conductivity, high hardness, and good chemical and thermal stability [1]. Previous works have indicated that the addition of SiC to HfB2 can improve the oxidation resistance and may improve the mechanical properties [2–5]. HfB2–SiC composites have been commonly produced by hot-pressing mechanically mixed powders. However, this method usually requires very high sintering temperature (above 1900 1C) and pressure (Z20 MPa), owing to the strong covalent bonding and low self-diffusion coefficients of HfB2 and SiC [6,7]. Reactive hot pressing shows a very good potential for sintering materials at a relative low temperature. HfB2–SiC composites have been successfully sintered using Hf, Si and B4C as starting powders [8,9]. Such a approach leads to composites having fine microstructure and isotropic properties. Spark plasma sintering (SPS) has been intensively used for preparing HfB2–SiC in recent years due to combination of unique properties such as fast heating/ n
Corresponding author. E-mail addresses: [email protected] (H. Wang), [email protected] (S.-H. Lee). http://dx.doi.org/10.1016/j.ceramint.2014.03.107 0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
cooling rate and high applicable pressure [10]. In the present work, HfB2–SiC composite was prepared by reactive SPS (R-SPS), using HfSi2, B4C and carbon as starting powders. The reaction mechanisms, densification process, mechanical properties and microstructures of the composites prepared by R-SPS are reported. 2. Experimental procedures Commercially available HfSi2 (1–2 μm, purity 499%; Alfa Aesar, MA, USA), B4C (purity 99%, particle size o 10 μm; Alfa Aesar, MA, USA) and carbon black powers (surface area 60–80 m2/g, purity 99.5%; Alfa Aesar, MA, USA) were used as starting materials. The starting powders were mixed with a stoichiometric composition according to the following reaction formula (1), targeting 28.6 wt% of SiC in the final HfB2–SiC composites. HfSi2 þ B4 C þ 3C ¼ 2HfB2 þ 4SiC
ð1Þ
The raw powders were mixed for 2 h using a shaker mill (Spex D8000, Spex CertPrep, Metuchen, NJ) equipped with two cylindrical WC-Co containers. WC-Co balls with 6.7 mm diameter were employed as milling media. The ball-to-powder weight ratio was 10. In order to prevent possible oxidation during mixing, the containers were sealed in a grove box with 95% nitrogenþ 5%
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H. Wang et al. / Ceramics International 40 (2014) 11009–11013
hydrogen atmosphere. After drying mixing, the obtained powders were dispersed by the shake mill for 5 min in wet conditions using ethanol as a dispersant. Then the slurry was homogenized again by high powder sonication for 30 min, and was dried by rotary evaporation. The power mixtures were loaded onto a graphite mold (10 mm inner diameter) lined with graphite foil and densified using SPS (Dr. Sinter 2020, Sumitomo Coal Mining Co., Tokyo, Japan) in a vacuum ( 6 Pa) at 1600 1C under a uniaxial pressure of 40 MPa (heating rate, 100 1C/min) for 10 min. A 12 ms-on and 2 ms-off pulse sequence was applied. Heating process was controlled by means of a monochromatic optical pyrometer that was focused on a hole at the side of the graphite mold. The sintering shrinkage of the specimen was analyzed in the manner of measuring the displacement of the lower electrode (resolution 0.01 mm) which was connected to a computer to log shrinkage curves. The bulk densities of sintered samples were determined using Archimedes’ method. The theoretical density of the HfB2–SiC composites calculated on the basis the Rule of Mixtures was 6.55 g/cm3. Young's modulus was determined through ultrasonic testing at 20–22 1C (ICP-MS. DRC 3000, Perkin/Elmer, USA), and the hardness and fracture toughness of the samples were obtained using a Vickers indenter (AVK-A, Akashi, Tokyo, Japan; loading condition: 1 kg, 15 s) according to the Anstis formula [11]: P Hv ¼ 1:854 2 ð2Þ d K IC ¼ 0:016
ðE=HvÞ1=2 P ðcÞ3=2
ð3Þ
where Hv means Vickers hardness (GPa); P is the applied force for indentation (N); d is the average diagonal length of indent (m); KIC is the fracture toughness (MPa m1/2); E is the Young's modulus (GPa); and c is the average crack length from the center of the indent to the crack tip (m). For each sample, 20 indentations were made and 40 diagonal lengths were measured to have a representative mean value of the hardness. The microstructure and phase fraction of sintered specimens were analyzed with scanning electron microscopy (SEM, JEOL JSM-6700F, Tokyo, Japan) and X-ray diffraction analysis (XRD, D/MAX 2200, Rigaku, Tokyo, Japan) using Cu Kα radiation. The microstructure and chemical composition of the specimens were further analyzed using a transmission electron microscopy (TEM, JEM 2100 F, JEOL, Tokyo, Japan) equipped with EDS (spot diameter: 1.0 nm). 3. Results and discussion 3.1. Densification behavior Fig. 1(a) shows shrinkage curves during sintering of the HfB2–SiC composite. Shrinkage was initiated (Tinitial) at 1250 1C. The peak shrinkage rate was analyzed through the differentiation of the shrinkage curves (Fig. 1(b)). It was found that the shrinkage of the sample progressed slowly in the first
Fig. 1. The displacement (a) and shrinkage rate (b) of the HfB2–SiC composites during reactive spark plasma sintering.
stage from 1250 1C to 1500 1C. Subsequently, rapid densification occurred between 1500 1C and 1600 1C. The maximum shrinkage rate was attained at 1550 1C. Finally, the shrinkage rate decreased continuously in the early stage of isothermal heating, and the shrinkage stopped after 100 s during isothermal heating. The shape of the curve corresponded to the mechanism of solid-state sintering [12]. In the stage of slow densification, the rearrangement of particles took place in the compact. The second stage was a sintering regime where accelerated densification occurred and the pore distribution transformed from open porosity to closed porosity [12]. The sintering behavior changed into the final stage of sintering during isothermal heating, where the closed porosity is eliminated. The final stage of sintering typically occurs when relative density becomes above 90% [12]. At the maximum shrinkage rate, the sintering mechanism changed from surface diffusion to grain-boundary diffusion (at 1550 1C) [13]. Finally, the HfB2–SiC composite became nearly fully dense (98.7%) during isothermal heating at 1600 1C for 10 min. Fig. 2 shows the XRD patterns of the starting powders and the sintered specimens. It was found that the size of the starting powders decreased to 10 nm (as calculated by the Scherrer formula) by high-energy ball milling. Consequently, the reaction (1) could easily proceed by decrease in the diffusion distance between ions. Also, high-energy ball milling promotes the densification of ultra-high-temperature ceramics by dramatically decreasing the particle size [14]. The peak intensity of HfB2 increased with the increasing sintering temperature. The reaction (1) was nearly completed, when the sintering temperature reached 1250 1C. The crystalline phases detected by X-ray diffraction in the composite were HfB2, SiC and a small amount of HfSi2 at 1250 1C. The result indicated that the densification of the composite proceeded mainly by solid-state sintering above 1250 1C and the exothermal reaction (1) did not strongly promote the densification. High-energy ball milling improved the reaction kinetics of the starting powders to form very fine HfB2 and SiC at a low temperature. The solid-state sintering of the HfB2–SiC composite was most probably promoted by the small particle size of the starting materials [15]. So the densification of the HfB2–SiC composite is caused both by the low temperature reaction between the
H. Wang et al. / Ceramics International 40 (2014) 11009–11013
Fig. 2. XRD patterns of the starting powder and sintered HfB2–SiC composite by SPS.
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Fig. 4 shows a typical HRTEM image and EDX data which were derived from the HfB2–SiC composites. An intergranular amorphous film was observed in the grain boundary between SiC grains (in Fig. 4(b)). An EDX spectrum of the grain boundary showed that only Si and C elements were present as shown in Fig. 4(d), which indicated that oxide phases such as SiO2 or B2O3 were not present at the grain boundary. The grain boundary between HfB2 and SiC was clean and the presence of an amorphous grain boundary was not observed (in Fig. 4(c)). An EDX spectrum of the SiC grain displayed that only Si and C elements were present within the grain (Fig. 3(f)). Moreover, the SEAD pattern corresponded to that of SiC, indicating the highly crystallized nature of the SiC. An EDX spectrum of the HfB2 grain showed the presence of Hf element. B was not detected due to the limit of the equipment (in Fig. 3(e)). The SEAD pattern corresponded to that of highly crystallized HfB2. The EDX analyses clearly indicated that the formation of solid solution between pure SiC and HfB2 did not occur during reactive SPS at 1600 1C. 3.3. Mechanical properties
Fig. 3. A typical bright-field transmission electron microscope image of the HfB2–SiC composite after sintering at 1600 1C for 10 min.
raw materials which was promoted by high-energy ball milling and by the fine particle size of HfB2 and SiC formed by the reaction (1).
Young's modulus (E) of the composites was 432 GPa, which was lower than the reported values (512–544 GPa) [1]. The distribution SiC grains boundary might result in a relatively low Young's modulus. The hardness of the composites was 20.471.5 GPa which was similar to the reported values of 17– 26 GPa [1]. The fracture toughness of the HfB2–SiC composite was 4.770.3 MPa m1/2, which was higher than the reported results of HfB2–SiC composite (3.9–4.1 MPa m1/2) [1]. In order to elucidate the toughening mechanisms, the propagation path of cracks and the fractured surface of the specimen were observed by SEM, as shown in Fig. 5. The fractured surface presented the mixture of trans- and inter-granular modes (Fig. 5(a)). Crack deflection and crack bridging were observed in the composites, which increased the resistance against crack propagation in the sample and consumed more energy for separation of the fractured surface (Fig. 5(b)). The HfB2 grains typically failed in a transgranular manner and cracks deflected at or near HfB2–SiC interfaces, leaving SiC particles in the wake of the advancing cracks. The result was caused by the residual stress which was accumulated at the interface of HfB2–SiC grains due to the mismatch of thermal and mechanical properties between SiC (CTE: 4.7 10 6 K 1) [16] and HfB2 (CTE: 6.3 10 6 K 1) [1]. Therefore, crack deflection and crack bridging were believed to be the primary toughening mechanism of HfB2–SiC composite.
3.2. Microstructure analysis Fig. 3 presents a bright-field TEM image of the HfB2–SiC composite. On the whole, the distribution of the in situ formed HfB2 and SiC phases was homogeneous. Most of the SiC grains had an equiaxed shape, and the grain size was smaller than 1 μm. However, a lot of the HfB2 grains grew and had an elongated-rod-like shape with an average aspect ratio of approx. 2. The average grain size of HfB2 was determined to be 1.5 μm. The small grain diameter is attributable to a lower sintering temperature and shorter holding time.
4. Conclusions HfB2–SiC composites could be in situ synthesized and densified by the reactive spark plasma sintering of HfSi2, B4C and C starting powders at a low temperature of 1600 1C. The reduction in particle size of the starting materials by high-energy ball milling and the significant enhancement of mass transfer during the reactive SPS process were found to be crucial factors for densification enhancement. The final microstructure was uniform and rather fine.
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Fig. 4. (a) A high-magnification, high-resolution transmission electron microscope image of grain boundary ((b) SiC grain boundary, (c) the boundary between SiC and HfB2) and energy dispersive X-ray spectroscopy spectra analysis ((d) derived from SiC grain boundary, (e) derived from HfB2 grain, (f) derived from SiC grain) of the HfB2–SiC composite.
Fig. 5. SEM images of the indentation crack path (a) and fracture surface (b) of the HfB2–SiC composite.
H. Wang et al. / Ceramics International 40 (2014) 11009–11013
The fracture toughness of the HfB2–SiC composite was 4.770.3 MPa m1/2, in combination with Young's modulus of 432 GPa and Vickers hardness of 20.471.5 GPa. The crack deflection and crack bridging were believed to be the primary toughening mechanisms of the HfB2–SiC composites. The residual stress in the composite can promote these two mechanisms. References [1] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347–1364. [2] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, Mechanical, thermal, and oxidation properties of refractory hafnium and zirconium compounds, J. Eur. Ceram. Soc. 19 (1999) 2405–2414. [3] X. Zhang, L. Weng, J. Han, S. Meng, W. Han, Preparation and thermal ablation behavior of HfB2–SiC-based ultra-high-temperature ceramics under severe heat conditions, Int. J. Appl. Ceram. Technol. 6 (2009) 134–144. [4] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics, J. Mater. Sci. 39 (2004) 5925–5937. [5] F. Monteverde, A. Bellosi, The resistance to oxidation of an HfB2-SiC composite, J. Eur. Ceram. Soc. 25 (2005) 1025–1031. [6] F. Monteverde, Ultra-high temperature HfB2-SiC ceramics consolidated by hot-pressing and spark plasma sintering, J. Alloys Compd. 428 (2007) 197–205.
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[7] J. Zou, G.J. Zhang, Y.M. Kan, T. Ohji, Pressureless sintering mechanisms and mechanical properties of hafnium diboride ceramics with presintering heat treatment, Scr. Mater. 62 (2010) 159–162. [8] F. Monteverde, Progress in the fabrication of ultra-high temperature ceramics: “in situ” synthesis, microstructure and properties of a reactive hot-pressed HfB2-SiC composite, Comp. Sci. Technol. 65 (2005) 1869–1879. [9] S.J. Lee, E.S. Kang, S.S. Baek, D.K. Kim, Reactive hot pressing and oxidation behavior of Hf-based ultra-high-temperature ceramics, Surf. Rev. Lett. 17 (2010) 215–221. [10] F. Monteverde, C. Melandri, S. Guicciardi, Microstructure and mechanical properties of an HfB2 þ 30 vol.% SiC composite consolidated by spark plasma sintering, Mater. Chem. Phys. 100 (2006) 513–519. [11] G.R. Anstis, P. Chantiklul, B.R. Lawn, D.B. Marshall, A critical evaluation of indentation techniques for measuring fracture toughness: I: direct crack measurements, J. Am. Ceram. Soc. 64 (1981) 533–538. [12] R.M. German, Sintering Theory and Practice, New York Wiley, 1996. [13] Z. Shen, M. Nygren, Microstructural prototyping of ceramics by kinetic engineering: applications of spark plasma sintering, Chem. Rec. 5 (2005) 173–184. [14] C.A. Galan, A.L. Ortiz, F. Guiberteau, High-Energy ball milling of ZrB2 in the presence of graphite, J. Am. Ceram. Soc. 93 (2012) 3072–3075. [15] V. Zamora, A.L. Ortiz, F. Guiberteau, M. Nygren, Crystal-size dependence of the spark-plasma-sintering kinetics of ZrB2 ultra-hightemperature ceramics, J. Eur. Ceram. Soc. 32 (2012) 271–276. [16] H.L. Wang, B.B. Fan, L. Feng, The fabrication and mechanical properties of SiC/ZrB2 laminated ceramic composite prepared by spark plasma sintering, Ceram. Int. 38 (2012) 5015–5022.
CERAMICS INTERNATIONAL
Ceramics International 40 (2014) 11009–11013 www.elsevier.com/locate/ceramint
HfB2–SiC composite prepared by reactive spark plasma sintering Hailong Wanga, Sea-Hoon Leeb,n, Lun Fengb a School of Materials Science and Engineering, Zhengzhou University, Zhengzhou 450001, China Division of Powder/Ceramics Research, Korea Institute of Materials Science, 531 Changwondaero, Changwon, Gyeongnam 641–831, Republic of Korea
b
Received 10 March 2014; received in revised form 19 March 2014; accepted 19 March 2014 Available online 27 March 2014
Abstract Dense HfB2–SiC composites were fabricated by reactive spark plasma sintering of a powder mixture composed of HfSi2, B4C and carbon. The relative density of the composites was 98.7% after reactive spark plasma sintering at 1600 1C under a pressure of 40 MPa for 10 min. Strong exothermic reaction occurred during the heating of the powder mixture due to the following reaction: 2HfSi2 þB4Cþ3C ¼2HfB2 þ 4SiC. The in situ formed HfB2 and SiC phases were homogeneously distributed. The average grain size of formed HfB2 and SiC was about 2 and 1 μm, respectively. The Vickers hardness and fracture toughness of HfB2–SiC composites were 20.47 1.5 GPa and 4.7 70.3 MPa m1/2, respectively. The present report proposes a method to densify HfB2-based composites at a low temperature. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: Hafnium diboride; Microstructure; Reactive spark plasma sintering
1. Introduction Hafnium diboride (HfB2) is a potential candidate for ultra-hightemperature applications in the aerospace industry because of its high melting point, high strength, high thermal conductivity, high hardness, and good chemical and thermal stability [1]. Previous works have indicated that the addition of SiC to HfB2 can improve the oxidation resistance and may improve the mechanical properties [2–5]. HfB2–SiC composites have been commonly produced by hot-pressing mechanically mixed powders. However, this method usually requires very high sintering temperature (above 1900 1C) and pressure (Z20 MPa), owing to the strong covalent bonding and low self-diffusion coefficients of HfB2 and SiC [6,7]. Reactive hot pressing shows a very good potential for sintering materials at a relative low temperature. HfB2–SiC composites have been successfully sintered using Hf, Si and B4C as starting powders [8,9]. Such a approach leads to composites having fine microstructure and isotropic properties. Spark plasma sintering (SPS) has been intensively used for preparing HfB2–SiC in recent years due to combination of unique properties such as fast heating/ n
Corresponding author. E-mail addresses: [email protected] (H. Wang), [email protected] (S.-H. Lee). http://dx.doi.org/10.1016/j.ceramint.2014.03.107 0272-8842/& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
cooling rate and high applicable pressure [10]. In the present work, HfB2–SiC composite was prepared by reactive SPS (R-SPS), using HfSi2, B4C and carbon as starting powders. The reaction mechanisms, densification process, mechanical properties and microstructures of the composites prepared by R-SPS are reported. 2. Experimental procedures Commercially available HfSi2 (1–2 μm, purity 499%; Alfa Aesar, MA, USA), B4C (purity 99%, particle size o 10 μm; Alfa Aesar, MA, USA) and carbon black powers (surface area 60–80 m2/g, purity 99.5%; Alfa Aesar, MA, USA) were used as starting materials. The starting powders were mixed with a stoichiometric composition according to the following reaction formula (1), targeting 28.6 wt% of SiC in the final HfB2–SiC composites. HfSi2 þ B4 C þ 3C ¼ 2HfB2 þ 4SiC
ð1Þ
The raw powders were mixed for 2 h using a shaker mill (Spex D8000, Spex CertPrep, Metuchen, NJ) equipped with two cylindrical WC-Co containers. WC-Co balls with 6.7 mm diameter were employed as milling media. The ball-to-powder weight ratio was 10. In order to prevent possible oxidation during mixing, the containers were sealed in a grove box with 95% nitrogenþ 5%
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H. Wang et al. / Ceramics International 40 (2014) 11009–11013
hydrogen atmosphere. After drying mixing, the obtained powders were dispersed by the shake mill for 5 min in wet conditions using ethanol as a dispersant. Then the slurry was homogenized again by high powder sonication for 30 min, and was dried by rotary evaporation. The power mixtures were loaded onto a graphite mold (10 mm inner diameter) lined with graphite foil and densified using SPS (Dr. Sinter 2020, Sumitomo Coal Mining Co., Tokyo, Japan) in a vacuum ( 6 Pa) at 1600 1C under a uniaxial pressure of 40 MPa (heating rate, 100 1C/min) for 10 min. A 12 ms-on and 2 ms-off pulse sequence was applied. Heating process was controlled by means of a monochromatic optical pyrometer that was focused on a hole at the side of the graphite mold. The sintering shrinkage of the specimen was analyzed in the manner of measuring the displacement of the lower electrode (resolution 0.01 mm) which was connected to a computer to log shrinkage curves. The bulk densities of sintered samples were determined using Archimedes’ method. The theoretical density of the HfB2–SiC composites calculated on the basis the Rule of Mixtures was 6.55 g/cm3. Young's modulus was determined through ultrasonic testing at 20–22 1C (ICP-MS. DRC 3000, Perkin/Elmer, USA), and the hardness and fracture toughness of the samples were obtained using a Vickers indenter (AVK-A, Akashi, Tokyo, Japan; loading condition: 1 kg, 15 s) according to the Anstis formula [11]: P Hv ¼ 1:854 2 ð2Þ d K IC ¼ 0:016
ðE=HvÞ1=2 P ðcÞ3=2
ð3Þ
where Hv means Vickers hardness (GPa); P is the applied force for indentation (N); d is the average diagonal length of indent (m); KIC is the fracture toughness (MPa m1/2); E is the Young's modulus (GPa); and c is the average crack length from the center of the indent to the crack tip (m). For each sample, 20 indentations were made and 40 diagonal lengths were measured to have a representative mean value of the hardness. The microstructure and phase fraction of sintered specimens were analyzed with scanning electron microscopy (SEM, JEOL JSM-6700F, Tokyo, Japan) and X-ray diffraction analysis (XRD, D/MAX 2200, Rigaku, Tokyo, Japan) using Cu Kα radiation. The microstructure and chemical composition of the specimens were further analyzed using a transmission electron microscopy (TEM, JEM 2100 F, JEOL, Tokyo, Japan) equipped with EDS (spot diameter: 1.0 nm). 3. Results and discussion 3.1. Densification behavior Fig. 1(a) shows shrinkage curves during sintering of the HfB2–SiC composite. Shrinkage was initiated (Tinitial) at 1250 1C. The peak shrinkage rate was analyzed through the differentiation of the shrinkage curves (Fig. 1(b)). It was found that the shrinkage of the sample progressed slowly in the first
Fig. 1. The displacement (a) and shrinkage rate (b) of the HfB2–SiC composites during reactive spark plasma sintering.
stage from 1250 1C to 1500 1C. Subsequently, rapid densification occurred between 1500 1C and 1600 1C. The maximum shrinkage rate was attained at 1550 1C. Finally, the shrinkage rate decreased continuously in the early stage of isothermal heating, and the shrinkage stopped after 100 s during isothermal heating. The shape of the curve corresponded to the mechanism of solid-state sintering [12]. In the stage of slow densification, the rearrangement of particles took place in the compact. The second stage was a sintering regime where accelerated densification occurred and the pore distribution transformed from open porosity to closed porosity [12]. The sintering behavior changed into the final stage of sintering during isothermal heating, where the closed porosity is eliminated. The final stage of sintering typically occurs when relative density becomes above 90% [12]. At the maximum shrinkage rate, the sintering mechanism changed from surface diffusion to grain-boundary diffusion (at 1550 1C) [13]. Finally, the HfB2–SiC composite became nearly fully dense (98.7%) during isothermal heating at 1600 1C for 10 min. Fig. 2 shows the XRD patterns of the starting powders and the sintered specimens. It was found that the size of the starting powders decreased to 10 nm (as calculated by the Scherrer formula) by high-energy ball milling. Consequently, the reaction (1) could easily proceed by decrease in the diffusion distance between ions. Also, high-energy ball milling promotes the densification of ultra-high-temperature ceramics by dramatically decreasing the particle size [14]. The peak intensity of HfB2 increased with the increasing sintering temperature. The reaction (1) was nearly completed, when the sintering temperature reached 1250 1C. The crystalline phases detected by X-ray diffraction in the composite were HfB2, SiC and a small amount of HfSi2 at 1250 1C. The result indicated that the densification of the composite proceeded mainly by solid-state sintering above 1250 1C and the exothermal reaction (1) did not strongly promote the densification. High-energy ball milling improved the reaction kinetics of the starting powders to form very fine HfB2 and SiC at a low temperature. The solid-state sintering of the HfB2–SiC composite was most probably promoted by the small particle size of the starting materials [15]. So the densification of the HfB2–SiC composite is caused both by the low temperature reaction between the
H. Wang et al. / Ceramics International 40 (2014) 11009–11013
Fig. 2. XRD patterns of the starting powder and sintered HfB2–SiC composite by SPS.
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Fig. 4 shows a typical HRTEM image and EDX data which were derived from the HfB2–SiC composites. An intergranular amorphous film was observed in the grain boundary between SiC grains (in Fig. 4(b)). An EDX spectrum of the grain boundary showed that only Si and C elements were present as shown in Fig. 4(d), which indicated that oxide phases such as SiO2 or B2O3 were not present at the grain boundary. The grain boundary between HfB2 and SiC was clean and the presence of an amorphous grain boundary was not observed (in Fig. 4(c)). An EDX spectrum of the SiC grain displayed that only Si and C elements were present within the grain (Fig. 3(f)). Moreover, the SEAD pattern corresponded to that of SiC, indicating the highly crystallized nature of the SiC. An EDX spectrum of the HfB2 grain showed the presence of Hf element. B was not detected due to the limit of the equipment (in Fig. 3(e)). The SEAD pattern corresponded to that of highly crystallized HfB2. The EDX analyses clearly indicated that the formation of solid solution between pure SiC and HfB2 did not occur during reactive SPS at 1600 1C. 3.3. Mechanical properties
Fig. 3. A typical bright-field transmission electron microscope image of the HfB2–SiC composite after sintering at 1600 1C for 10 min.
raw materials which was promoted by high-energy ball milling and by the fine particle size of HfB2 and SiC formed by the reaction (1).
Young's modulus (E) of the composites was 432 GPa, which was lower than the reported values (512–544 GPa) [1]. The distribution SiC grains boundary might result in a relatively low Young's modulus. The hardness of the composites was 20.471.5 GPa which was similar to the reported values of 17– 26 GPa [1]. The fracture toughness of the HfB2–SiC composite was 4.770.3 MPa m1/2, which was higher than the reported results of HfB2–SiC composite (3.9–4.1 MPa m1/2) [1]. In order to elucidate the toughening mechanisms, the propagation path of cracks and the fractured surface of the specimen were observed by SEM, as shown in Fig. 5. The fractured surface presented the mixture of trans- and inter-granular modes (Fig. 5(a)). Crack deflection and crack bridging were observed in the composites, which increased the resistance against crack propagation in the sample and consumed more energy for separation of the fractured surface (Fig. 5(b)). The HfB2 grains typically failed in a transgranular manner and cracks deflected at or near HfB2–SiC interfaces, leaving SiC particles in the wake of the advancing cracks. The result was caused by the residual stress which was accumulated at the interface of HfB2–SiC grains due to the mismatch of thermal and mechanical properties between SiC (CTE: 4.7 10 6 K 1) [16] and HfB2 (CTE: 6.3 10 6 K 1) [1]. Therefore, crack deflection and crack bridging were believed to be the primary toughening mechanism of HfB2–SiC composite.
3.2. Microstructure analysis Fig. 3 presents a bright-field TEM image of the HfB2–SiC composite. On the whole, the distribution of the in situ formed HfB2 and SiC phases was homogeneous. Most of the SiC grains had an equiaxed shape, and the grain size was smaller than 1 μm. However, a lot of the HfB2 grains grew and had an elongated-rod-like shape with an average aspect ratio of approx. 2. The average grain size of HfB2 was determined to be 1.5 μm. The small grain diameter is attributable to a lower sintering temperature and shorter holding time.
4. Conclusions HfB2–SiC composites could be in situ synthesized and densified by the reactive spark plasma sintering of HfSi2, B4C and C starting powders at a low temperature of 1600 1C. The reduction in particle size of the starting materials by high-energy ball milling and the significant enhancement of mass transfer during the reactive SPS process were found to be crucial factors for densification enhancement. The final microstructure was uniform and rather fine.
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Fig. 4. (a) A high-magnification, high-resolution transmission electron microscope image of grain boundary ((b) SiC grain boundary, (c) the boundary between SiC and HfB2) and energy dispersive X-ray spectroscopy spectra analysis ((d) derived from SiC grain boundary, (e) derived from HfB2 grain, (f) derived from SiC grain) of the HfB2–SiC composite.
Fig. 5. SEM images of the indentation crack path (a) and fracture surface (b) of the HfB2–SiC composite.
H. Wang et al. / Ceramics International 40 (2014) 11009–11013
The fracture toughness of the HfB2–SiC composite was 4.770.3 MPa m1/2, in combination with Young's modulus of 432 GPa and Vickers hardness of 20.471.5 GPa. The crack deflection and crack bridging were believed to be the primary toughening mechanisms of the HfB2–SiC composites. The residual stress in the composite can promote these two mechanisms. References [1] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347–1364. [2] M.M. Opeka, I.G. Talmy, E.J. Wuchina, J.A. Zaykoski, S.J. Causey, Mechanical, thermal, and oxidation properties of refractory hafnium and zirconium compounds, J. Eur. Ceram. Soc. 19 (1999) 2405–2414. [3] X. Zhang, L. Weng, J. Han, S. Meng, W. Han, Preparation and thermal ablation behavior of HfB2–SiC-based ultra-high-temperature ceramics under severe heat conditions, Int. J. Appl. Ceram. Technol. 6 (2009) 134–144. [4] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, Processing, properties and arc jet oxidation of hafnium diboride/silicon carbide ultra high temperature ceramics, J. Mater. Sci. 39 (2004) 5925–5937. [5] F. Monteverde, A. Bellosi, The resistance to oxidation of an HfB2-SiC composite, J. Eur. Ceram. Soc. 25 (2005) 1025–1031. [6] F. Monteverde, Ultra-high temperature HfB2-SiC ceramics consolidated by hot-pressing and spark plasma sintering, J. Alloys Compd. 428 (2007) 197–205.
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