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Densification, microstructure evolution and mechanical properties of WC doped HfB2–SiC ceramics
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ScienceDirect Journal of the European Ceramic Society 35 (2015) 2707–2714
Densification, microstructure evolution and mechanical properties of WC doped HfB2–SiC ceramics Ji-Xuan Liu a , Guo-Jun Zhang a,∗ , Fang-Fang Xu a , Wen-Wen Wu b , Hai-Tao Liu a , Yoshio Sakka b , Toshiyuki Nishimura b , Tohru S. Suzuki b , De-Wei Ni a , Ji Zou a b
a Shanghai Institute of Ceramics, CAS, Shanghai 200050, China Materials Processing Unit, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Ibaraki, Japan
Received 13 July 2014; received in revised form 4 April 2015; accepted 7 April 2015 Available online 19 April 2015
Abstract Despite of good oxidation resistance and ablation resistance, challenge to densify and tendency to loss high temperature (H.T.) strength have limited its potential applications in aerospace for HfB2 –SiC ceramic. In this work, dense HfB2 –SiC ceramic with improved H.T. flexural strength was prepared using WC as sintering aid. Pure HfB2 –SiC ceramic exhibited creep deformation at 1600 ◦ C with measured strength value of 389 ± 82 MPa, whereas WC doped sample showed much higher flexural strength of 658 ± 69 MPa with linear elastic behavior prior to fracture. In addition to clean grain boundary, the addition of 5 vol.%WC into HfB2 –SiC ceramic promoted the formation of fine (Hf,W)B2 , (Hf,W)C and WB precipitates. This combinative effect is responsible for its excellent H.T. mechanical property of WC doped HfB2 –SiC ceramic. © 2015 Elsevier Ltd. All rights reserved. Keywords: UHTCs; Hafnium diboride; High-temperature properties; Sintering; Ceramic
1. Introduction Belong to the family of UHTCs of Group IV metal carbides, borides and nitrides, Hafnium diboride (HfB2 ) and Zirconium diboride (ZrB2 ) have unique combinations of thermophysical properties, including high melting point, high modulus and hardness, high thermal conductivity and high resistance to chemical attack [1–5]. Compared to ZrB2 , oxidation resistance of HfB2 is better [1]. As a result, HfB2 -based ceramics have been considered as promising candidates for ultra-high temperature applications, such as propulsion systems, rocket nozzles, sharp leading edges and nose cones. It is believed that some of those may be able to operate over 3000 ◦ C. Early studies showed that oxidation resistance and mechanical properties of HfB2 based ceramics can be improved significantly by silicon carbide addition [1,5]. Therefore, HfB2 –SiC system has attracted great attention in recent years.
∗
Corresponding author. Tel.: +86 21 52411080; fax: +86 21 52413122. E-mail address: [email protected] (G.-J. Zhang).
http://dx.doi.org/10.1016/j.jeurceramsoc.2015.04.009 0955-2219/© 2015 Elsevier Ltd. All rights reserved.
Due to its strong covalent bonding and low self-diffusion coefficient, HfB2 -based ceramic is difficult to densify [1,6]. In addition, impurities such as HfO2 and B2 O3 , often located on surfaces of HfB2 powders, inhibiting densification. In order to enhance the densification, hot pressing or processing aids for pressureless sintering started with fine source powders are common routines of practice. HfB2 –SiC ceramic with relative density of 99.2% has been hot pressed at 2000 ◦ C using fine HfB2 powder [7]. Several silicides and carbides, such as HfSi2 , MoSi2 , TaSi2 , B4 C and WC, were selected as sintering additives to enhance densification of HfB2 via either liquid phase sintering or oxygen removal mechanisms [6,8–13]. For example, sample closing to full density with average HfB2 grain size of 4 m was obtained from the addition of 5 vol.%HfSi2 doped HfB2 ceramic. The enhancement of mass transportation and acceleration of densification was ascribed to the formation of Hf–Si liquid phase around 1600 ◦ C during hot pressing [8]. Wang et al. [11], prepared HfB2 –SiC ceramic with relative density of 98.9% and average HfB2 grain size of 3 m by hot pressing at 1850 ◦ C using 10 vol.%B4 C as sintering additive. WC has been selected as sintering additive for boride ceramics
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to accelerate densification and improve oxidation resistance of HfB2 –SiC ceramics [12–15]. To ensure stable load-carrying capacity of component under service condition, HfB2 –SiC ceramics should have excellent flexural strength at H.T. [16]. However, the flexural strength of HfB2 –SiC ceramics decreased rapidly above 1000 ◦ C. The four-point flexural strength of hot pressed HfB2 –15 vol.%SiC ceramics decreased from 460 MPa to 400 MPa from room temperature (R.T.) to 1450 ◦ C [17]. How to maintain their H.T. strength is vital for real applications. In previous works, ZrB2 –SiC ceramics with improved H.T. flexural strength was prepared using high purity ZrB2 powder as raw material and WC as sintering aid [18,19]. For pure ZrB2 –SiC ceramic, threepoint flexural strength of 460 MPa was registered at 1600 ◦ C, while that of WC doped sample exhibited enhanced value of 675 MPa. Removal of oxide impurities from grain boundaries was considered to be the main reason for such high flexural strength [19]. Thus, it is expected that hot pressed HfB2 –SiC ceramics with similar material properties can be obtained from high purity HfB2 powder or the addition of WC. In the present study, high purity fine HfB2 powder was selected as raw material, the densifications of hot pressed HfB2 –20 vol.%SiC with or without WC addition were analyzed. The effect of WC addition on microstructure evolution and mechanical properties as function of temperature was investigated and discussed.
2. Experimental procedure 2.1. Powder processing and sintering HfB2 powder (purity > 99%, oxygen content 0.15 wt.%, C < 0.04 wt.%, mean particle size 1 m) synthesized from hafnia and B4 C [20], commercial ␣-SiC powder (purity > 98.5%, oxygen content 1.04 wt.%, mean particle size 450 nm, Changle Xinyuan Carborundum Micropowder Co., LTD, Changle, China) and WC powder (purity > 99%, mean particle size 0.8 m, Zhuzhou Cemented Carbide Group Co., LTD, Zhuzhou, China) were used as raw materials. Samples with compositions of 80 vol.%HfB2 –20 vol.%SiC (abbreviated as HS) and extra 5 vol.%WC addition in HS (abbreviated as HSW) were prepared. The starting powders with or without WC additive were mixed in ethanol by rolling ball mill at speed of 30 rpm for 24 h, using Si3 N4 balls as medium. Subsequently, the slurries were dried in rotary evaporator at 80 ◦ C. The dried powders were sieved through 200 mesh screen. Hot pressing was conducted in H.T. graphite resistant furnace (ZT-60-22Y, Chenhua Furnaces Inc., Shanghai, China). The mixed powder compacts in graphite die were heated from R.T. to 1650 ◦ C at a rate of 10 ◦ C/min and dwell for 0.5 h in vacuum. After the hold at 1650 ◦ C, the furnace atmosphere was changed to flowing argon (purity 99.99%) at pressure of ∼105 Pa. The powder compact was continually heated to 2000 ◦ C at 15 ◦ C/min under uniaxial pressure of 30 MPa. After dwelling at 2000 ◦ C for 1 h, the applied pressure was removed. The obtained sintered samples were sectioned, ground and polished to 0.5 m
finish using either diamond paste or silica gel for successive characterization. 2.2. Sample characterization Densities of sintered samples were measured using the Archimedes method. Oxygen contents of samples were measured by oxygen analyzer (TC600, LECO, St Joseph, MI, USA). Sintered samples were crushed into small pieces in steel mold and then ground manually using an agate mortar and pestle. Finally, powders sieved through 60 mesh screen were used for oxygen contents measurement. Phase composition was analyzed by X-ray diffraction (XRD, D/max 2250 V, Rigaku, Tokyo, Japan) using Cu K␣ radiation and using Si as internal standard. Microstructures were observed using scanning electron microscope (SEM, JSM-6700F, JEOL Ltd., Tokyo, Japan) and high resolution transmission electron microscope (HR-TEM, JEM-2010, JEOL Ltd., Tokyo, Japan) equipped with energy dispersive X-ray spectroscopy (EDS). Volume fractions and average grain size of phases were determined through image analysis of SEM micrographs. Vickers hardness (HV ) and fracture toughness (KIc ) were determined using Vickers indentation (2100B Hardness Tester, Instron, Norwood, MA, USA) under load of 5 kg and dwell time of 10 s [21]. All values of HV and KIc were the average of five measurements. Flexural strength tests were conducted on Instron-5500-universal testing machine (Instron, Norwood, MA, USA) in air. H.T. flexural strength was measured on H.T. bending test machine (Model 4505, Instron, Norwood, MA, USA) in flowing argon. In order to prevent oxidation of the samples during H.T. flexural strength measurements, the testing chamber was first evacuated down to 10−3 Pa, followed by backfilling high purity argon (>99.999%) at 100 mL/min. The furnace was raised to desired temperature at rate of 30 ◦ C/min. To reach thermal equilibrium for the specimen, load was applied after soaking at testing temperatures for 20 min. Flexural strength of four-point bending at various temperatures (25, 1300, 1600 ◦ C) was tested on 2.5 mm × 2 mm × 30 mm bars using 10 mm and 20 mm as inner and outer spans, respectively. The crosshead speed was 0.5 mm/min [9,16,18,19]. All bars were chamfered on their edges and polished with 0.5 m diamond paste before bending test. The reported flexural strengths are the average of five specimens at R.T. and three specimens at H.T., respectively. 3. Results and discussion 3.1. Oxygen content Oxygen contamination (including B2 O3 , ZrO2 , HfO2 ) is one of the major factors inhibiting densification of boride-based UHTCs. Besides, oxide impurities located at grain boundary are harmful to H.T. mechanical properties of boride-based UHTCs. Oxide impurities softened at H.T. and decreased H.T. flexural strength of ZrB2 -based ceramics [16–18]. In order to accelerate densification and improve H.T. mechanical properties of HfB2 -based ceramics, oxide impurities must be removed prior to densification.
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Fig. 1. The variations of gas pressure of furnace chamber (a) and the displacements of rods (b) during hot-pressing HS and HSW ceramics.
Owning to high purity of starting HfB2 powder, residual oxygen content in sintered HS ceramic was as low as 0.16 wt.%. The residual oxygen content in sintered HSW ceramic was only 0.02 wt.%. These results showed that WC addition was helpful for the removal of oxide impurities in HfB2 –SiC ceramics. Oxide impurities on the surface of SiC and HfB2 particles most likely take the form of SiO2 , B2 O3 and HfO2 . SiO2 impurity can be removed by reaction between SiC and SiO2 at temperatures above 1700 ◦ C. This processing is often called as self-cleaning of SiC [22]. B2 O3 will be vaporized significantly at temperatures above 1200 ◦ C. Unlike B2 O3 , HfO2 with melting point of 2810 ◦ C has very low vapor pressure in the sintering temperature range. To remove HfO2 , some sintering additives are desired. Our previous work revealed that WC could react with HfO2 and remove the impurity via reactions (1) and (2) during pressureless sintering at H.T. [12]: HfO2 + 3WC → HfC + 3W + 2CO(g) (T > 1260 ◦ C, PCO = 10 Pa, ΔG < 0)
(1)
HfO2 + 6WC + 3HfB2 → 4HfC + 6WB + 2CO(g) (T > 1160 ◦ C, PCO = 10 Pa, ΔG < 0)
(2)
Thermodynamic calculation shows that both reactions (1) and (2) become favorable above ∼1200 ◦ C at CO(g) partial pressure of 10 Pa. In the present experiment, gas pressure of furnace chamber began to rise at ∼1200 ◦ C during sintering of HSW ceramic in vacuum (see Fig. 1a), this revealed above mentioned two gas evolution reactions (Eqs. (1) and (2)) have indeed occurred. These reactions effectively reduced HfO2 impurity and leaded to low oxygen content of sintered HSW ceramic. 3.2. Densification The relative densities of hot pressed HS and HSW ceramics were 98.6% and 99%, respectively. The results indicate the starting powders have good sinterability. Fig. 1b shows
the displacements of rods during hot-pressing of HS and HSW ceramics. The more rapid densification of HSW is as expected from previous data concerning densification of diborides containing WC additives [13–15]. 3.3. Phase composition and microstructure of sintered ceramics Fig. 2 shows typical TEM pictures and related EDS results of HS and HSW ceramics. There is one thin amorphous oxide layer on grain boundary of HfB2 in HS ceramic (see Fig. 2a and b). In contrast, the grain boundary of HSW is clean (see Fig. 2c and d), which is attributed to the oxygen removal effect of WC additive. Related EDS results revealed that both HfC and HfB2 grains in HSW ceramics contained W element (see Fig. 2e and f). Thus, (Hf,W)B2 and (Hf,W)C solid solutions formed during sintering of HSW ceramics. XRD patterns of the as-sintered ceramics indicated HfB2 and SiC were the main phases (see Fig. 3a). Besides, HfC and WB peaks were only found in XRD pattern of HSW ceramics. The HfB2 and HfC peaks in XRD pattern of HSW ceramics were shifted to higher 2θ values due to the formation of (Hf,W)B2 and (Hf,W)C solid solutions with lattice distortion (see Fig. 3b) [12,13]. In addition, the lattice distortion could decrease activation energy of sintering and thus improve densification of HSW ceramics. Fig. 4 shows BSE morphologies of polished surfaces and fracture surfaces of HS and HSW ceramics. According to XRD analysis, the gray, black, light gray and white phases on the polished surface of HSW ceramic are (Hf,W)B2 , SiC, (Hf,W)C and WB, respectively (see Fig. 4c). Volume fractions (determined by image analysis) of (Hf,W)B2 , SiC, (Hf,W)C and WB in HSW ceramic are about 78%, 19.3%, 0.7% and 2%, respectively. Average grain sizes of HfB2 and SiC in HS ceramics are about 2.5 m and 2 m, respectively. In contrast, WC addition significantly refined the microstructure of HSW ceramics. Average grain sizes of (Hf,W)B2 and SiC in HSW ceramics are only 1.5 m. While average grain sizes of WB and (Hf,W)C phase in HSW ceramics are only 0.8 m. Although the polished surfaces of HS and HSW ceramics showed uniform microstructure, several SiC clusters with sizes about 30 m
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Fig. 2. Typical TEM pictures and related EDS results of HS (a), (b) and HSW (c)–(f) ceramics.
Fig. 3. XRD patterns of HS and HSW ceramics (a) and enlarged image of the shifted diffraction peaks (b).
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Fig. 4. BSE micrographs of polished surfaces (a, c) and fracture surfaces (b, d) of HS and HSW ceramics.
were found on related fracture surfaces (see Fig. 4b and d). Large SiC clusters can cause high levels of residual stress and stress concentration, and harmful to mechanical properties of boride–SiC composite ceramics [23]. In addition, microcracks existed in HS ceramic (see Figs. 2a and 4a), possibly due to thermal expansion mismatch between HfB2 (6.3 × 10−6 K−1 at 298 K) and SiC (3.3 × 10−6 K−1 at 298 K) grains [7,24]. 3.4. Mechanical properties Vickers hardness of HS and HSW ceramics were measured to be 19.5 ± 0.8 GPa and 22.3 ± 1.5 GPa, respectively. Higher hardness of HSW ceramic was due to hard WB phase (Hv = 30.1 GPa). Fracture toughness of HS and HSW ceramics were determined as 3.95 ± 0.4 MPa·m1/2 and 3.76 ± 0.7 MPa·m1/2 , respectively. Fig. 5 shows the four-point flexural strength of HS and HSW as a function of temperature. R.T. flexural strengths of HS and HSW were 526 ± 86 MPa and 544 ± 135 MPa, respectively. The flexural strength of ceramics is inversely proportional to critical flaw present in the microstructure, as described by Griffith–Irwin equation [25,26]. σ=
KIc Y × a−1/2
and stress concentration had been identified as the critical flaw in ZrB2 –SiC ceramics [23]. Thus, it is reasonable to conclude that SiC clusters are also the strength-limiting flaws in HS and HSW ceramics at R.T. The measured flexural strengths of HS ceramics at 1300 ◦ C and 1600 ◦ C were 419 ± 13 MPa and 389 ± 82 MPa, respectively. Flexural strengths of HS ceramics were higher than the reported values of hot pressed HfB2 –20 vol.%SiC ceramics (see Fig. 5) [3,27]. Even though high purity HfB2 powders were used as raw material here, we still could not prevent strength degradation of HS ceramics from happening. However, HSW ceramics exhibited high strength values at elevated temperatures. Flexural strength of HSW ceramics at 1300 ◦ C was 563 ± 61 MPa.
(3)
here σ is flexural strength, Y is geometric constant, and a is critical flaw size. Y can be 1.28 for boride–SiC composite ceramic. Based on Eq. (3), the calculated critical flaw sizes in HS and HSW ceramics are 34 m and 29 m, respectively, which are similar to the size of SiC clusters (30 m) in sintered ceramics. In previous work, SiC inclusions causing large residual stress
Fig. 5. Four-point flexural strength of HS and HSW ceramics vs. temperature.
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Fig. 6. Load-displacement curves of HS and HSW ceramics during bending test (a) and appearances of fractured bars (b).
Especially, flexural strength of HSW ceramics increased to 658 ± 69 MPa at 1600 ◦ C. These results showed WC addition can significantly improve H.T. flexural strength of HfB2 –SiC ceramics.
(see Figs. 4b and 7e and f). This was attributed to the softening of oxide impurities on grain boundaries of HS. In contrast, HSW ceramics with less oxide impurities had more stable grain boundaries and thus showed transgranular fracture up to 1600 ◦ C (see Figs. 4d and 7g and h).
3.5. Deformation during H.T. bending test The load-displacement curves of HS and HSW ceramics during bending test at various temperatures are shown in Fig. 6a. All sintered ceramics exhibited elastic deformation till fractured at below 1300 ◦ C. Pure HfB2 –SiC ceramic exhibited creep deformation at 1600 ◦ C, whereas WC doped sample showed linear elastic behavior prior to fracture. Appearances of fractured HS and HSW bars after bending test are in accordance with the loaddisplacement curves (see Fig. 6b). Only HS bars softened during bending test at 1600 ◦ C. In addition, previous works showed that the measured strengths of UHTCs at elevated temperatures are loading rate dependent [23,28]. To get an accurate value of H.T. strength, the crosshead speed should be varied with temperature such that the failure of tested samples occurred in a linear elastic fashion [23,28,29]. Thus, further investigations on H.T. flexural strengths of HfB2 –SiC ceramics tested under different loading rates are needed. 3.6. Microstructure evolution during H.T. bending test Fig. 7 shows SEM micrographs of free surfaces on tensile side of HS and HSW ceramics after H.T. bending test and the related fracture surfaces. Some oxide impurities could evaporate from grain boundaries during H.T. bending test, forming grain boundary grooves on the surfaces of HS ceramics (white arrows in Fig. 7a and b). In contrast, no grain boundary groove existed on the surfaces of HSW ceramics (see Fig. 7c and d). The thermal stability of grain boundaries in HSW ceramics is much higher than those in HS ceramics. HS ceramics showed transgranular fracture at R.T., however, it exhibited intergranular fracture at H.T. temperatures
3.7. Effect of oxide impurities, SiC clusters and residual stress on flexural strength Previous works showed that SiC clusters, residual stress and softening of the grain boundary phase strongly influenced flexural strength of dense, fine grained ZrB2 –SiC ceramics [18,23,24,30]. These factors also influenced mechanical properties of HS and HSW ceramics here in this work. R.T. flexural strength of HS and HSW ceramics was mainly controlled by SiC clusters, as discussed in section 3.4. But H.T. flexural strength of sintered ceramics was strongly affected by other factors. The amorphous grain boundary phase in HS ceramics softened at temperatures above ∼1000 ◦ C, which resulted in degradation of flexural strength and creep deformation of HS ceramics at elevated temperatures [18]. On the other hand, WC addition effectively cleaned the grain boundaries of HSW ceramics. Thus, flexural strength of HSW ceramics didn’t decrease at elevated temperatures. Especially, flexural strength of HSW ceramics increased at 1600 ◦ C, which was attributed to residual stress relaxation. Residual stress in boride–SiC composite ceramics often begins to accumulate at ∼1400 ◦ C upon cooling from processing temperature [23,24]. It could induce microcracks and consequently decreased flexural strength of HSW ceramic below 1400 ◦ C. Oppositely, residual stress relaxed through diffusional mechanism at temperatures above 1400 ◦ C [24]. Thus, fine grained HSW ceramics with clean grain boundaries showed improved flexural strength at 1600 ◦ C. On the basis of above discussions, it is concluded that the elimination of SiC agglomerates, the removal of oxide impurities and relaxation of residual stress can improve H.T. flexural strength of fine grained HfB2 –SiC ceramics.
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Fig. 7. SEM micrographs of free surfaces on tensile side (a)–(d) of HS and HSW ceramics after H.T. bending test and the related fracture surfaces (e)–(h).
4. Conclusions Dense HfB2 –20 vol.%SiC ceramics with and without WC addition were prepared by hot pressing at 2000 ◦ C. The measured flexural strengths of pure HfB2 –SiC ceramics at R.T., 1300 ◦ C and 1600 ◦ C were 526 ± 86 MPa, 419 ± 13 MPa and 389 ± 82 MPa, respectively. Flexural strengths of WC doped samples measured at R.T., 1300 ◦ C and 1600 ◦ C were 544 ± 135 MPa, 563 ± 61 MPa and 658 ± 69 MPa, respectively. Pure HfB2 –SiC ceramic exhibited creep deformation at 1600 ◦ C, whereas WC doped sample showed linear elastic behavior prior
to fracture. In addition to clean grain boundary, the addition of 5 vol.%WC into HfB2 –SiC ceramic promoted the formation of fine (Hf,W)B2 , (Hf,W)C and WB precipitates. This combinative effect is responsible for its excellent H.T. mechanical property of WC doped HfB2 –SiC ceramic. Acknowledgements Financial supports from the National Natural Science Foundation of China (No. 51272266), the Science and Technology Commission of Shanghai (No. 15ZR1445200) and the State
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Key Laboratory of High Performance Ceramics and Superfine Microstructure are gratefully acknowledged. References [1] Fahrenholtz WG, Wuchina EJ, Lee WE, Zhou YC. Ultra-High Temperature Ceramics: Materials for Extreme Environment Applications. Hoboken, NJ: Published by John Wiley & Sons, Inc.; 2014. [2] Wuchina E, Opila E, Opeka MM, Fahrenholtz W, Talmy I. UHTCs: ultrahigh temperature ceramic materials for extreme environment applications. Interface 2007;16:30–6. [3] Fahrenholtz WG, Hilmas GE. Refractory diborides of zirconium and hafnium. J Am Ceram Soc 2007;90:1347–64. [4] Wuchina E, Opeka M, Causey S, Buesking K, Spain J, Cull A, et al. Designing for ultrahigh-temperature applications: the mechanical and thermal properties of HfB2 , HfCx , HfNx and ␣Hf(N). J Mater Sci 2004;39:5939–49. [5] Monteverde F, Bellosi A. Microstructure and properties of an HfB2 –SiC composite for ultra high temperature applications. Adv Eng Mater 2004;6:331–6. [6] Shaklee HJB, Fahrenholtz WG, Hilmas GE. Densification behavior and microstructure evolution of hot-pressed HfB2 . J Am Ceram Soc 2011;94:49–58. [7] Ni DW, Zhang GJ, Kan YM, Wang PL. Hot pressed HfB2 and HfB2 –20 vol.%SiC ceramics based on HfB2 powder synthesized by borothermal reduction of HfO2 . Int J Appl Ceram Technol 2010;7:830–6. [8] Monteverde F. Hot pressing of hafnium diboride aided by different sinter additives. J Mater Sci 2008;43:1002–7. [9] Sciti D, Silvestroni L, Bellosi A. Fabrication and properties of HfB2 –MoSi2 composites produced by hot pressing and spark plasma sintering. J Mater Res 2006;21:1460–6. [10] Hu CF, Sakka Y, Jang BK, Tanaka H, Nishimura T, Guo SQ, et al. Microstructure and properties of ZrB2 –SiC and HfB2 –SiC composites fabricated by spark plasma sintering (SPS) using TaSi2 as sintering aid. J Ceram Soc Jpn 2010;118:997–1001. [11] Wang L, Zhang XH, Han JC, Han WB, Hong CQ. The effect of B4 C on the microstructure and thermo-mechanical properties of HfB2 -based ceramics. J Alloys Compd 2009;473(1–2):314–8. [12] Ni DW, Liu JX, Zhang GJ. Pressureless sintering of HfB2 –SiC ceramics doped with WC. J Eur Ceram Soc 2012;32:3627–35. [13] Carney CM, Parthasarathy TA, Cinibulk MK. Oxidation resistance of hafnium diboride ceramics with additions of silicon carbide and tungsten boride or tungsten carbide. J Am Ceram Soc 2011;94:2600–7.
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ScienceDirect Journal of the European Ceramic Society 35 (2015) 2707–2714
Densification, microstructure evolution and mechanical properties of WC doped HfB2–SiC ceramics Ji-Xuan Liu a , Guo-Jun Zhang a,∗ , Fang-Fang Xu a , Wen-Wen Wu b , Hai-Tao Liu a , Yoshio Sakka b , Toshiyuki Nishimura b , Tohru S. Suzuki b , De-Wei Ni a , Ji Zou a b
a Shanghai Institute of Ceramics, CAS, Shanghai 200050, China Materials Processing Unit, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Ibaraki, Japan
Received 13 July 2014; received in revised form 4 April 2015; accepted 7 April 2015 Available online 19 April 2015
Abstract Despite of good oxidation resistance and ablation resistance, challenge to densify and tendency to loss high temperature (H.T.) strength have limited its potential applications in aerospace for HfB2 –SiC ceramic. In this work, dense HfB2 –SiC ceramic with improved H.T. flexural strength was prepared using WC as sintering aid. Pure HfB2 –SiC ceramic exhibited creep deformation at 1600 ◦ C with measured strength value of 389 ± 82 MPa, whereas WC doped sample showed much higher flexural strength of 658 ± 69 MPa with linear elastic behavior prior to fracture. In addition to clean grain boundary, the addition of 5 vol.%WC into HfB2 –SiC ceramic promoted the formation of fine (Hf,W)B2 , (Hf,W)C and WB precipitates. This combinative effect is responsible for its excellent H.T. mechanical property of WC doped HfB2 –SiC ceramic. © 2015 Elsevier Ltd. All rights reserved. Keywords: UHTCs; Hafnium diboride; High-temperature properties; Sintering; Ceramic
1. Introduction Belong to the family of UHTCs of Group IV metal carbides, borides and nitrides, Hafnium diboride (HfB2 ) and Zirconium diboride (ZrB2 ) have unique combinations of thermophysical properties, including high melting point, high modulus and hardness, high thermal conductivity and high resistance to chemical attack [1–5]. Compared to ZrB2 , oxidation resistance of HfB2 is better [1]. As a result, HfB2 -based ceramics have been considered as promising candidates for ultra-high temperature applications, such as propulsion systems, rocket nozzles, sharp leading edges and nose cones. It is believed that some of those may be able to operate over 3000 ◦ C. Early studies showed that oxidation resistance and mechanical properties of HfB2 based ceramics can be improved significantly by silicon carbide addition [1,5]. Therefore, HfB2 –SiC system has attracted great attention in recent years.
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Corresponding author. Tel.: +86 21 52411080; fax: +86 21 52413122. E-mail address: [email protected] (G.-J. Zhang).
http://dx.doi.org/10.1016/j.jeurceramsoc.2015.04.009 0955-2219/© 2015 Elsevier Ltd. All rights reserved.
Due to its strong covalent bonding and low self-diffusion coefficient, HfB2 -based ceramic is difficult to densify [1,6]. In addition, impurities such as HfO2 and B2 O3 , often located on surfaces of HfB2 powders, inhibiting densification. In order to enhance the densification, hot pressing or processing aids for pressureless sintering started with fine source powders are common routines of practice. HfB2 –SiC ceramic with relative density of 99.2% has been hot pressed at 2000 ◦ C using fine HfB2 powder [7]. Several silicides and carbides, such as HfSi2 , MoSi2 , TaSi2 , B4 C and WC, were selected as sintering additives to enhance densification of HfB2 via either liquid phase sintering or oxygen removal mechanisms [6,8–13]. For example, sample closing to full density with average HfB2 grain size of 4 m was obtained from the addition of 5 vol.%HfSi2 doped HfB2 ceramic. The enhancement of mass transportation and acceleration of densification was ascribed to the formation of Hf–Si liquid phase around 1600 ◦ C during hot pressing [8]. Wang et al. [11], prepared HfB2 –SiC ceramic with relative density of 98.9% and average HfB2 grain size of 3 m by hot pressing at 1850 ◦ C using 10 vol.%B4 C as sintering additive. WC has been selected as sintering additive for boride ceramics
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to accelerate densification and improve oxidation resistance of HfB2 –SiC ceramics [12–15]. To ensure stable load-carrying capacity of component under service condition, HfB2 –SiC ceramics should have excellent flexural strength at H.T. [16]. However, the flexural strength of HfB2 –SiC ceramics decreased rapidly above 1000 ◦ C. The four-point flexural strength of hot pressed HfB2 –15 vol.%SiC ceramics decreased from 460 MPa to 400 MPa from room temperature (R.T.) to 1450 ◦ C [17]. How to maintain their H.T. strength is vital for real applications. In previous works, ZrB2 –SiC ceramics with improved H.T. flexural strength was prepared using high purity ZrB2 powder as raw material and WC as sintering aid [18,19]. For pure ZrB2 –SiC ceramic, threepoint flexural strength of 460 MPa was registered at 1600 ◦ C, while that of WC doped sample exhibited enhanced value of 675 MPa. Removal of oxide impurities from grain boundaries was considered to be the main reason for such high flexural strength [19]. Thus, it is expected that hot pressed HfB2 –SiC ceramics with similar material properties can be obtained from high purity HfB2 powder or the addition of WC. In the present study, high purity fine HfB2 powder was selected as raw material, the densifications of hot pressed HfB2 –20 vol.%SiC with or without WC addition were analyzed. The effect of WC addition on microstructure evolution and mechanical properties as function of temperature was investigated and discussed.
2. Experimental procedure 2.1. Powder processing and sintering HfB2 powder (purity > 99%, oxygen content 0.15 wt.%, C < 0.04 wt.%, mean particle size 1 m) synthesized from hafnia and B4 C [20], commercial ␣-SiC powder (purity > 98.5%, oxygen content 1.04 wt.%, mean particle size 450 nm, Changle Xinyuan Carborundum Micropowder Co., LTD, Changle, China) and WC powder (purity > 99%, mean particle size 0.8 m, Zhuzhou Cemented Carbide Group Co., LTD, Zhuzhou, China) were used as raw materials. Samples with compositions of 80 vol.%HfB2 –20 vol.%SiC (abbreviated as HS) and extra 5 vol.%WC addition in HS (abbreviated as HSW) were prepared. The starting powders with or without WC additive were mixed in ethanol by rolling ball mill at speed of 30 rpm for 24 h, using Si3 N4 balls as medium. Subsequently, the slurries were dried in rotary evaporator at 80 ◦ C. The dried powders were sieved through 200 mesh screen. Hot pressing was conducted in H.T. graphite resistant furnace (ZT-60-22Y, Chenhua Furnaces Inc., Shanghai, China). The mixed powder compacts in graphite die were heated from R.T. to 1650 ◦ C at a rate of 10 ◦ C/min and dwell for 0.5 h in vacuum. After the hold at 1650 ◦ C, the furnace atmosphere was changed to flowing argon (purity 99.99%) at pressure of ∼105 Pa. The powder compact was continually heated to 2000 ◦ C at 15 ◦ C/min under uniaxial pressure of 30 MPa. After dwelling at 2000 ◦ C for 1 h, the applied pressure was removed. The obtained sintered samples were sectioned, ground and polished to 0.5 m
finish using either diamond paste or silica gel for successive characterization. 2.2. Sample characterization Densities of sintered samples were measured using the Archimedes method. Oxygen contents of samples were measured by oxygen analyzer (TC600, LECO, St Joseph, MI, USA). Sintered samples were crushed into small pieces in steel mold and then ground manually using an agate mortar and pestle. Finally, powders sieved through 60 mesh screen were used for oxygen contents measurement. Phase composition was analyzed by X-ray diffraction (XRD, D/max 2250 V, Rigaku, Tokyo, Japan) using Cu K␣ radiation and using Si as internal standard. Microstructures were observed using scanning electron microscope (SEM, JSM-6700F, JEOL Ltd., Tokyo, Japan) and high resolution transmission electron microscope (HR-TEM, JEM-2010, JEOL Ltd., Tokyo, Japan) equipped with energy dispersive X-ray spectroscopy (EDS). Volume fractions and average grain size of phases were determined through image analysis of SEM micrographs. Vickers hardness (HV ) and fracture toughness (KIc ) were determined using Vickers indentation (2100B Hardness Tester, Instron, Norwood, MA, USA) under load of 5 kg and dwell time of 10 s [21]. All values of HV and KIc were the average of five measurements. Flexural strength tests were conducted on Instron-5500-universal testing machine (Instron, Norwood, MA, USA) in air. H.T. flexural strength was measured on H.T. bending test machine (Model 4505, Instron, Norwood, MA, USA) in flowing argon. In order to prevent oxidation of the samples during H.T. flexural strength measurements, the testing chamber was first evacuated down to 10−3 Pa, followed by backfilling high purity argon (>99.999%) at 100 mL/min. The furnace was raised to desired temperature at rate of 30 ◦ C/min. To reach thermal equilibrium for the specimen, load was applied after soaking at testing temperatures for 20 min. Flexural strength of four-point bending at various temperatures (25, 1300, 1600 ◦ C) was tested on 2.5 mm × 2 mm × 30 mm bars using 10 mm and 20 mm as inner and outer spans, respectively. The crosshead speed was 0.5 mm/min [9,16,18,19]. All bars were chamfered on their edges and polished with 0.5 m diamond paste before bending test. The reported flexural strengths are the average of five specimens at R.T. and three specimens at H.T., respectively. 3. Results and discussion 3.1. Oxygen content Oxygen contamination (including B2 O3 , ZrO2 , HfO2 ) is one of the major factors inhibiting densification of boride-based UHTCs. Besides, oxide impurities located at grain boundary are harmful to H.T. mechanical properties of boride-based UHTCs. Oxide impurities softened at H.T. and decreased H.T. flexural strength of ZrB2 -based ceramics [16–18]. In order to accelerate densification and improve H.T. mechanical properties of HfB2 -based ceramics, oxide impurities must be removed prior to densification.
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Fig. 1. The variations of gas pressure of furnace chamber (a) and the displacements of rods (b) during hot-pressing HS and HSW ceramics.
Owning to high purity of starting HfB2 powder, residual oxygen content in sintered HS ceramic was as low as 0.16 wt.%. The residual oxygen content in sintered HSW ceramic was only 0.02 wt.%. These results showed that WC addition was helpful for the removal of oxide impurities in HfB2 –SiC ceramics. Oxide impurities on the surface of SiC and HfB2 particles most likely take the form of SiO2 , B2 O3 and HfO2 . SiO2 impurity can be removed by reaction between SiC and SiO2 at temperatures above 1700 ◦ C. This processing is often called as self-cleaning of SiC [22]. B2 O3 will be vaporized significantly at temperatures above 1200 ◦ C. Unlike B2 O3 , HfO2 with melting point of 2810 ◦ C has very low vapor pressure in the sintering temperature range. To remove HfO2 , some sintering additives are desired. Our previous work revealed that WC could react with HfO2 and remove the impurity via reactions (1) and (2) during pressureless sintering at H.T. [12]: HfO2 + 3WC → HfC + 3W + 2CO(g) (T > 1260 ◦ C, PCO = 10 Pa, ΔG < 0)
(1)
HfO2 + 6WC + 3HfB2 → 4HfC + 6WB + 2CO(g) (T > 1160 ◦ C, PCO = 10 Pa, ΔG < 0)
(2)
Thermodynamic calculation shows that both reactions (1) and (2) become favorable above ∼1200 ◦ C at CO(g) partial pressure of 10 Pa. In the present experiment, gas pressure of furnace chamber began to rise at ∼1200 ◦ C during sintering of HSW ceramic in vacuum (see Fig. 1a), this revealed above mentioned two gas evolution reactions (Eqs. (1) and (2)) have indeed occurred. These reactions effectively reduced HfO2 impurity and leaded to low oxygen content of sintered HSW ceramic. 3.2. Densification The relative densities of hot pressed HS and HSW ceramics were 98.6% and 99%, respectively. The results indicate the starting powders have good sinterability. Fig. 1b shows
the displacements of rods during hot-pressing of HS and HSW ceramics. The more rapid densification of HSW is as expected from previous data concerning densification of diborides containing WC additives [13–15]. 3.3. Phase composition and microstructure of sintered ceramics Fig. 2 shows typical TEM pictures and related EDS results of HS and HSW ceramics. There is one thin amorphous oxide layer on grain boundary of HfB2 in HS ceramic (see Fig. 2a and b). In contrast, the grain boundary of HSW is clean (see Fig. 2c and d), which is attributed to the oxygen removal effect of WC additive. Related EDS results revealed that both HfC and HfB2 grains in HSW ceramics contained W element (see Fig. 2e and f). Thus, (Hf,W)B2 and (Hf,W)C solid solutions formed during sintering of HSW ceramics. XRD patterns of the as-sintered ceramics indicated HfB2 and SiC were the main phases (see Fig. 3a). Besides, HfC and WB peaks were only found in XRD pattern of HSW ceramics. The HfB2 and HfC peaks in XRD pattern of HSW ceramics were shifted to higher 2θ values due to the formation of (Hf,W)B2 and (Hf,W)C solid solutions with lattice distortion (see Fig. 3b) [12,13]. In addition, the lattice distortion could decrease activation energy of sintering and thus improve densification of HSW ceramics. Fig. 4 shows BSE morphologies of polished surfaces and fracture surfaces of HS and HSW ceramics. According to XRD analysis, the gray, black, light gray and white phases on the polished surface of HSW ceramic are (Hf,W)B2 , SiC, (Hf,W)C and WB, respectively (see Fig. 4c). Volume fractions (determined by image analysis) of (Hf,W)B2 , SiC, (Hf,W)C and WB in HSW ceramic are about 78%, 19.3%, 0.7% and 2%, respectively. Average grain sizes of HfB2 and SiC in HS ceramics are about 2.5 m and 2 m, respectively. In contrast, WC addition significantly refined the microstructure of HSW ceramics. Average grain sizes of (Hf,W)B2 and SiC in HSW ceramics are only 1.5 m. While average grain sizes of WB and (Hf,W)C phase in HSW ceramics are only 0.8 m. Although the polished surfaces of HS and HSW ceramics showed uniform microstructure, several SiC clusters with sizes about 30 m
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Fig. 2. Typical TEM pictures and related EDS results of HS (a), (b) and HSW (c)–(f) ceramics.
Fig. 3. XRD patterns of HS and HSW ceramics (a) and enlarged image of the shifted diffraction peaks (b).
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Fig. 4. BSE micrographs of polished surfaces (a, c) and fracture surfaces (b, d) of HS and HSW ceramics.
were found on related fracture surfaces (see Fig. 4b and d). Large SiC clusters can cause high levels of residual stress and stress concentration, and harmful to mechanical properties of boride–SiC composite ceramics [23]. In addition, microcracks existed in HS ceramic (see Figs. 2a and 4a), possibly due to thermal expansion mismatch between HfB2 (6.3 × 10−6 K−1 at 298 K) and SiC (3.3 × 10−6 K−1 at 298 K) grains [7,24]. 3.4. Mechanical properties Vickers hardness of HS and HSW ceramics were measured to be 19.5 ± 0.8 GPa and 22.3 ± 1.5 GPa, respectively. Higher hardness of HSW ceramic was due to hard WB phase (Hv = 30.1 GPa). Fracture toughness of HS and HSW ceramics were determined as 3.95 ± 0.4 MPa·m1/2 and 3.76 ± 0.7 MPa·m1/2 , respectively. Fig. 5 shows the four-point flexural strength of HS and HSW as a function of temperature. R.T. flexural strengths of HS and HSW were 526 ± 86 MPa and 544 ± 135 MPa, respectively. The flexural strength of ceramics is inversely proportional to critical flaw present in the microstructure, as described by Griffith–Irwin equation [25,26]. σ=
KIc Y × a−1/2
and stress concentration had been identified as the critical flaw in ZrB2 –SiC ceramics [23]. Thus, it is reasonable to conclude that SiC clusters are also the strength-limiting flaws in HS and HSW ceramics at R.T. The measured flexural strengths of HS ceramics at 1300 ◦ C and 1600 ◦ C were 419 ± 13 MPa and 389 ± 82 MPa, respectively. Flexural strengths of HS ceramics were higher than the reported values of hot pressed HfB2 –20 vol.%SiC ceramics (see Fig. 5) [3,27]. Even though high purity HfB2 powders were used as raw material here, we still could not prevent strength degradation of HS ceramics from happening. However, HSW ceramics exhibited high strength values at elevated temperatures. Flexural strength of HSW ceramics at 1300 ◦ C was 563 ± 61 MPa.
(3)
here σ is flexural strength, Y is geometric constant, and a is critical flaw size. Y can be 1.28 for boride–SiC composite ceramic. Based on Eq. (3), the calculated critical flaw sizes in HS and HSW ceramics are 34 m and 29 m, respectively, which are similar to the size of SiC clusters (30 m) in sintered ceramics. In previous work, SiC inclusions causing large residual stress
Fig. 5. Four-point flexural strength of HS and HSW ceramics vs. temperature.
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Fig. 6. Load-displacement curves of HS and HSW ceramics during bending test (a) and appearances of fractured bars (b).
Especially, flexural strength of HSW ceramics increased to 658 ± 69 MPa at 1600 ◦ C. These results showed WC addition can significantly improve H.T. flexural strength of HfB2 –SiC ceramics.
(see Figs. 4b and 7e and f). This was attributed to the softening of oxide impurities on grain boundaries of HS. In contrast, HSW ceramics with less oxide impurities had more stable grain boundaries and thus showed transgranular fracture up to 1600 ◦ C (see Figs. 4d and 7g and h).
3.5. Deformation during H.T. bending test The load-displacement curves of HS and HSW ceramics during bending test at various temperatures are shown in Fig. 6a. All sintered ceramics exhibited elastic deformation till fractured at below 1300 ◦ C. Pure HfB2 –SiC ceramic exhibited creep deformation at 1600 ◦ C, whereas WC doped sample showed linear elastic behavior prior to fracture. Appearances of fractured HS and HSW bars after bending test are in accordance with the loaddisplacement curves (see Fig. 6b). Only HS bars softened during bending test at 1600 ◦ C. In addition, previous works showed that the measured strengths of UHTCs at elevated temperatures are loading rate dependent [23,28]. To get an accurate value of H.T. strength, the crosshead speed should be varied with temperature such that the failure of tested samples occurred in a linear elastic fashion [23,28,29]. Thus, further investigations on H.T. flexural strengths of HfB2 –SiC ceramics tested under different loading rates are needed. 3.6. Microstructure evolution during H.T. bending test Fig. 7 shows SEM micrographs of free surfaces on tensile side of HS and HSW ceramics after H.T. bending test and the related fracture surfaces. Some oxide impurities could evaporate from grain boundaries during H.T. bending test, forming grain boundary grooves on the surfaces of HS ceramics (white arrows in Fig. 7a and b). In contrast, no grain boundary groove existed on the surfaces of HSW ceramics (see Fig. 7c and d). The thermal stability of grain boundaries in HSW ceramics is much higher than those in HS ceramics. HS ceramics showed transgranular fracture at R.T., however, it exhibited intergranular fracture at H.T. temperatures
3.7. Effect of oxide impurities, SiC clusters and residual stress on flexural strength Previous works showed that SiC clusters, residual stress and softening of the grain boundary phase strongly influenced flexural strength of dense, fine grained ZrB2 –SiC ceramics [18,23,24,30]. These factors also influenced mechanical properties of HS and HSW ceramics here in this work. R.T. flexural strength of HS and HSW ceramics was mainly controlled by SiC clusters, as discussed in section 3.4. But H.T. flexural strength of sintered ceramics was strongly affected by other factors. The amorphous grain boundary phase in HS ceramics softened at temperatures above ∼1000 ◦ C, which resulted in degradation of flexural strength and creep deformation of HS ceramics at elevated temperatures [18]. On the other hand, WC addition effectively cleaned the grain boundaries of HSW ceramics. Thus, flexural strength of HSW ceramics didn’t decrease at elevated temperatures. Especially, flexural strength of HSW ceramics increased at 1600 ◦ C, which was attributed to residual stress relaxation. Residual stress in boride–SiC composite ceramics often begins to accumulate at ∼1400 ◦ C upon cooling from processing temperature [23,24]. It could induce microcracks and consequently decreased flexural strength of HSW ceramic below 1400 ◦ C. Oppositely, residual stress relaxed through diffusional mechanism at temperatures above 1400 ◦ C [24]. Thus, fine grained HSW ceramics with clean grain boundaries showed improved flexural strength at 1600 ◦ C. On the basis of above discussions, it is concluded that the elimination of SiC agglomerates, the removal of oxide impurities and relaxation of residual stress can improve H.T. flexural strength of fine grained HfB2 –SiC ceramics.
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Fig. 7. SEM micrographs of free surfaces on tensile side (a)–(d) of HS and HSW ceramics after H.T. bending test and the related fracture surfaces (e)–(h).
4. Conclusions Dense HfB2 –20 vol.%SiC ceramics with and without WC addition were prepared by hot pressing at 2000 ◦ C. The measured flexural strengths of pure HfB2 –SiC ceramics at R.T., 1300 ◦ C and 1600 ◦ C were 526 ± 86 MPa, 419 ± 13 MPa and 389 ± 82 MPa, respectively. Flexural strengths of WC doped samples measured at R.T., 1300 ◦ C and 1600 ◦ C were 544 ± 135 MPa, 563 ± 61 MPa and 658 ± 69 MPa, respectively. Pure HfB2 –SiC ceramic exhibited creep deformation at 1600 ◦ C, whereas WC doped sample showed linear elastic behavior prior
to fracture. In addition to clean grain boundary, the addition of 5 vol.%WC into HfB2 –SiC ceramic promoted the formation of fine (Hf,W)B2 , (Hf,W)C and WB precipitates. This combinative effect is responsible for its excellent H.T. mechanical property of WC doped HfB2 –SiC ceramic. Acknowledgements Financial supports from the National Natural Science Foundation of China (No. 51272266), the Science and Technology Commission of Shanghai (No. 15ZR1445200) and the State
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