Journal of the European Ceramic Society xxx (xxxx) xxx–xxx Contents lists available at ScienceDirect Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc Short communication Harmonized toughening and strengthening in pressurelessly reactivesintered Ta0.8Hf0.2C-SiC composite ⁎ Buhao Zhanga,d, Jie Yina,b, , Yihua Huanga, Jian Chena, Xuejian Liua, Zhengren Huanga,c, ⁎ a State Key Laboratory of High Performance Ceramics and Superﬁne Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China b Bureau of Major R&D Programs, Chinese Academy of Sciences, Beijing 100864, China c Ningbo Institute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo 315201, China d University of Chinese Academy of Sciences, Beijing 100049, China A R T I C LE I N FO A B S T R A C T Keywords: Ta0.8Hf0.2C-SiC composite Pressureless in-situ densiﬁcation Phase transformation Toughen Strengthen Ta0.8Hf0.2C-27 vol%SiC (99.0% in relative density) composite was toughened and strengthened via pressurelessly in-situ reactive sintering process. HfC and β-SiC particles were formed after reaction of HfSi2 and carbon black at 1650 °C. Ta0.8Hf0.2C was obtained from solid solutioning of HfC and commercial TaC. The β-α phase transformation of SiC proceeded below 2200 °C. High aspect ratio, platelet-like α-SiC grains formed and interconnected as interlocking structures. Toughness and ﬂexural strength values of 5.4 ± 1.2 MPa m1/2 and 443 ± 22 MPa were measured respectively. The toughening mechanisms by highly directional growth of discontinuous α-SiC grains were crack branching, bridging and deﬂection behaviors. 1. Introduction Ta1-xHfxC, belonging to the family of ultra-high temperature ceramics (UHTCs), has attracted considerable attention due to high melting point, superior thermal and chemical stabilities . The solid solution formation not only contributes to densiﬁcation by decreasing the diffusion activation energy across grain boundaries , but may also improve mechanical properties than their monolithic carbide counterparts: relationship between hardness and x values of Ta1-xHfxC (x = 0, 0.2, 0.5 and 0.8) was reported to vary as an open-down quadratic function, which could be attributed to solid solution strengthening effect . Notably, Ta0.8Hf0.2C was proved to have the highest melting point 3942 ± 82 °C among the Ta1-xHfxC solid solutions . Ta0.8Hf0.2C is the candidate material for thermal protection systems (TPS) as sharp leading edges, nose caps and ﬂight control components of aerospace vehicles that can be exposed to extreme environments with temperatures exceeding 2000 °C by dissociated air operate at hypersonic speeds. Owing to its extremely strong covalent bonding, Ta0.8Hf0.2C is difﬁcult to densify. Advanced techniques including hot pressing, spark plasma sintering  and hot isostatic pressing  have been attempted to enhance the sintering driving force. Pressureless sintering (PLS) is beneﬁcial to fabricate cost-eﬀective and near-net-shape ceramic components. Researches on the ﬁeld-less densiﬁcation techniques of Ta0.8Hf0.2C were rather scarce. Quite recently, we reported pressureless densiﬁcation of Ta0.8Hf0.2C-based composites with ﬁne mechanical/ thermal properties . The critical application of Ta0.8Hf0.2C-based composites as thermal protection systems (TPS) is restricted by the poor thermal shock resistance, including thermal shock fracture and damage resistance. The maximum thermal shock fracture parameter (R) is one of crucial importance to impede the fracture inside bulk ceramics . Thermal shock damage resistance is considered on the basis of toughness (KIC), with the pre-existing cracks after fracture initiation, when the maximum stress intensity factor Kmax reaches the material toughness KIC . Therefore, UHTCs need to improve their damage tolerance via increasing both strength and toughness for TPS application . To meet such demand, incorporating a hard secondary phase to strengthen, together with further extrinsic toughening , such as crack deﬂection and bridging is an eﬀective strategy. Investigations are needed for the optimization of Ta0.8Hf0.2C-SiC mechanical performance, typically insitu strengthening and toughening. SiC, introduced in-situ inside the (Ta,Hf)C bulk, were tailored to be platelet-like with high aspect ratio of 15.6 after soaking 2 h at 2200 °C. Homogeneous HfC and SiC formed after reaction between HfSi2 and ⁎ Corresponding authors at: State Key Laboratory of High Performance Ceramics and Superﬁne Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China. E-mail addresses: email@example.com (B. Zhang), firstname.lastname@example.org (J. Yin), email@example.com (Z. Huang). https://doi.org/10.1016/j.jeurceramsoc.2018.08.021 Received 24 May 2018; Received in revised form 16 August 2018; Accepted 17 August 2018 0955-2219/ © 2018 Published by Elsevier Ltd. Please cite this article as: Zhang, B., Journal of the European Ceramic Society (2018), https://doi.org/10.1016/j.jeurceramsoc.2018.08.021 Journal of the European Ceramic Society xxx (xxxx) xxx–xxx B. Zhang et al. Fig. 1. (a) Thermodynamics analysis of the reactions for TaC-HfSi2-carbon black system; XRD patterns of (b) HfSi2 and carbon black reaction and (c) HfSi2 and carbon black reaction in TaC matrix from [33°, 43°] at diﬀerent temperatures. hardness was measured by the indentation technique (Model 300, Tukon, Canton, MA, USA) using a load of 3 kg and dwell time of 10 s. The indentation fracture resistance (KIC) was calculated on the basis of the equation as reported by Evans and Charles . carbon black at 1650 °C. Ta0.8Hf0.2C was generated via inter-diﬀusion among raw TaC and in-situ HfC. The phase and microstructure evolution of SiC proceeded during the high-temperature densiﬁcation progress. The in-situ toughening mechanism was investigated. KIC = 0.16(c/a)−1.5∙H∙a1/2 2. Experimental procedure Where KIC is fracture toughness, a is the half average length of the diagonal of the Vickers indentations (μm), c is the average length of the cracks obtained in the tips of the Vickers indentations (μm), and H is the Vickers hardness. TaC and HfSi2 powders (Haotian Nano Technology Co., Ltd, Shanghai, China) had a purity of 99.9% and an average particle size of 1 μm. The median size of carbon black (Anyang Delong Chemical Co., Ltd., Henan, China) was ∼200 nm. Phenolic resin (P. R.) (Shanghai QiNan Adhesive Material Factory, Shanghai, China) with a carbon yield of 50 wt% was used as binder. The oxygen contents of TaC and HfSi2 were analyzed using infrared method after fusion under an inert gas atmosphere (TC-600C, LECO Instrument LD, MI, USA). TaC, HfSi2 and carbon contents were precisely calculated for obtaining Ta0.8Hf0.2C27 vol% SiC in stoichiometric and volume ratio based on the reaction below : HfSi 2 + 3C = HfC + 2SiC ΔG = −167385−7.456T 3. Results and discussion To obtain desired properties of (Ta,Hf)C-SiC composite, the chemical reactions in TaC-HfSi2-C system were investigated. Oxide impurities were coated on the surfaces of raw powders, eg: Ta2O5 on TaC (oxygen contents: 0.59 wt%), and HfO2 and SiO2 on HfSi2 (oxygen contents: 1.73 wt%). The complexity of the chemical reactions in TaCHfSi2-C system increases due to the surface oxide impurities. Carbon was an eﬀective reducing agent. Following chemical reactions should be considered [13,14]: (1) Where ΔG : Gibbs’ free energy (J/mol), and T: Temperature (K). The starting powders were blended with 4 wt% phenolic resin in ethanol, and ball-milled for 4 h using a planetary mill (QM-3SP4, Nanjing NanDa Instrument plant, Nanjing, China) with WC balls (5 mm in diameter) and Teﬂon-coated tanks (Φ 100 mm × 120 mm) at a speed of 400 rpm. Wear debris contamination was ≤0.2 wt % after milling. The slurries were dried, crushed, sieved and pressed uniaxially in a steel die at 45 MPa (6 mm × 8 mm × 48 mm), followed by cold isostatic pressing at 280 MPa for 300 s. The green compacts were pressurelessly sintered at 1650 °C and 2200 °C for 0, 1 and 2 h in argon (purity 99.999% with O2 ≤ 1.5 ppm), at a heating rate of 10 °C min−1 in a graphite resistance furnace (Zhuzhou Norbert High Temperature Instrument Ltd. Co., China). Phase analysis was performed by X-ray diﬀraction (XRD; Ultima IV diﬀractometer, Rigaku, Tokyo, Japan) with Cu Kα radiation (λ = 1.5406 Å). The microstructure and crack introduced by indentation in hardness test on polished surface were analyzed by scanning electron microscopy (SEM; Magellan 400, FEI, Hillsboro, American) equipped with EBSD (INCA SERIES, Oxford Instrument, UK). The electron backscatter diﬀraction pattern (EBSP) was acquired at an angle of 60° and an acceleration voltage of 15 kV, and Aztec software was used to clarify the phase. In transmission electron microscopy (TEM; Tecnai G2 F20, FEI Co., Hillsboro, USA) observation, high-resolution transmission electron microscopy (HRTEM) was conducted. The bulk density was measured by the Archimedes method. The theoretical density was determined based on the rule of mixtures (densities of Ta0.8Hf0.2C and SiC are 14.05 and 3.21 g cm−3 respectively). Three-point bending strength (3 mm × 4 mm × 36 mm) was measured by a universal tester (Instron-1195, Instron, Canton, MA, USA) using a 30 mm span and a cross-head speed of 0.5 mm min−1. Vickers Ta2O5 + 7C → 2TaC+5CO(g) (2) HfO2 + 3C → HfC + 2CO(g) (3) SiO2 + 3C → SiC + 2CO(g) (4) 2SiO2 + SiC → 3SiO(g) + CO(g) (5) The Gibbs free energy (ΔG) of each reaction was calculated by thermodynamic simulation software (HSC Chemistry 6.1), and they were shown in Fig.1a. To complete the reactions above, adequate amount of carbon black was added to react with oxide impurities. The thermal decomposition of P. R. occurred before densiﬁcation initiation, during which all P. R. will be removed from the green body at 1000 °C. 0.9 wt% of pyrolytic carbon from P. R. removal was added into the total carbon contents precisely for obtaining Ta0.8Hf0.2C-27 vol% SiC. Reactions (2)–(4) occurs since 1600 °C, while Reaction (5) starts at above 1900 °C as shown in Fig. 1a. Besides, since SiO2 impurity could not be eliminated during the sintering of monolithic SiC without the presence of carbon additive, Reactions (2)–(4) were considered to occur for removing the surface impurities. Reaction (1) was taken out and investigated individually. The XRD patterns of HfSi2-carbon black sintered at 900 °C (P. R. removal) and 1650 °C (reaction (1) initiation) were shown in Fig.1b. Peaks of mixed HfSi2 and carbon black were consistent with the HfSi2 (PDF#38-1373) at 900 °C. Well-deﬁned HfC (PDF#39-1491) and β-SiC (PDF#29-1129, 3C-SiC) peaks were detected after the sintering at and above 1650 °C together with the complete disappearance of HfSi2 peaks, which indicated that reaction (1) was completed. After the temperature reached 2200 °C, existence of α-SiC phases in Fig.1b (PDF#29-1127, PDF#492 Journal of the European Ceramic Society xxx (xxxx) xxx–xxx B. Zhang et al. Fig. 2. EBSD phase-map of Ta0.8Hf0.2C-SiC composite; EBSP of Ta0.8Hf0.2C, 4H-SiC and 6H-SiC particles in the composite. Fig. 3(e), new α-SiC nuclei were assumed to grow inside the β-SiC grain, resulting in β/α hybrid SiC grains at above 1650 °C. Strain at the β/α interface might inﬂuence the orientation growth of ﬁnal elongated α-SiC grains during the temperature-increasing stage . The anisotropic α-SiC grains were randomly oriented and uniformly distributed in the Ta0.8Hf0.2C matrix. The relative density of Ta0.8Hf0.2C-27 vol%SiC increased from 87.2% to 99.0% after soaking at 2200 °C, as shown in Table 1. Soaking at 2200 °C led to limited grain growth of Ta0.8Hf0.2C matrix from 11.2 μm to 13.1 μm. Interestingly, the aspect ratio of the elongated α-SiC grains increased remarkably from 6.0 to 15.2 after 2 h’s dwelling at 2200 °C. Highly directional growth (platelet-like) of α-SiC grains was generated. α-SiC grains interconnected with each other as interlocking structures in Fig. 3d. This was speculated to improve the damage tolerance signiﬁcantly . The thermal expansion coeﬃcient mismatch between the Ta0.8Hf0.2C and SiC  during sintering was presumably to be responsible: localized stresses generated at the grain boundary could guide the directional grain growth of the hexagonal α-SiC during the dwelling at 2200 °C. HRTEM result of the sample sintered at 2200 °C for 2 h revealed the Ta0.8Hf0.2C and α-SiC grain boundary was clean, without any evidence of inter-diﬀusion between Ta0.8Hf0.2C and α-SiC in Fig. 3f. Additionally, grain coarsening of the matrix phase (Ta0.8Hf0.2C) proceeded mainly during the 2200 °C soaking period based our previous research . As a result, localized stresses at the grain boundary of Ta0.8Hf0.2C and α-SiC may be further accumulated and therefore lead to evolution of the highaspect-ratio grains [7,21] and detailed study is currently under way. The mechanical properties of Ta0.8Hf0.2C-based composites sintered at 2200 °C for 2 h were shown in Table 2. The Vickers hardness of Ta0.8Hf0.2C-27 vol%SiC composite was 16.0 ± 0.1 GPa, inheriting well from the monolithic Ta0.8Hf0.2C ceramic. The ﬂexural strength of the composite increased to 443 ± 22 MPa, 23.5% higher than the dense Ta0.8Hf0.2C ceramic. Toughness improvement in the composite was good. Toughness values of up to 5.4 ± 1.2 MPa m1/2 was measured, 45.9% higher than the SiC free sample, 3.7 ± 0.1 MPa m1/2. Crack path was marked by the Vickers indentation technique as shown in Fig. 4. Platelet-like α-SiC grains obviously deﬂected the crack propagation pathway into more 1428, 4 H/6H-SiC) implied a β-α SiC phase transformation. (Ta,Hf)C solid solution here was formed via mutual diﬀusion between commercial cubic TaC and in-situ generated cubic HfC phase (PDF#39-1491 by XRD veriﬁcation). Limited carbon vacancy  could stabilize in the cubic carbide lattices under carbon-deﬁcient condition, however, the carbon-rich environment during sintering (both carbon black and phenolic resin were included in the raw recipe) was presumed to contribute to the removal of carbon vacancy in our (Ta,Hf)C solid solution lattice. The lattice parameter of 2200 °C sintered (Ta,Hf)C was calculated from the XRD to be 4.4858 ± 0.0003 Å, which followed the Vegard’s Law (4.4899 Å) well. A Ta0.8Hf0.2C-based composite was formed with the in-situ introduction of α-SiC at 2200 °C. (Ta,Hf)C solid solution formation was investigated thoroughly in our previous work . Ta atoms can ‘dissolve’ more eﬃciently into the HfC lattice on account of higher diﬀusion activation energy of HfC . The main (Ta,Hf)C peak shifted from (34.63°, 1650 °C) to (34.70°, 2200 °C) in Fig. 1c. EBSD phase map and EBSP results of polished surface of Ta0.8Hf0.2C27 vol% SiC sintered at 2200 °C were exhibited in Fig. 2. EBSP results showed the presences of Ta0.8Hf0.2C, (Cubic, 225 Space Group, m-3 m Laue group) and α-SiC (4 H/6H-SiC, Hexagonal, 186 Space Group, 6/ mmm Laue group). The distribution of α-SiC within the Ta0.8Hf0.2C matrix was homogeneous and no β-SiC was observed according to the EBSD map. This conﬁrmed that the phase transformation of β→α SiC was completed after sintering at 2200 °C. By applying the internal standard method for quantitative analysis of XRD at 2200 °C , the composition of Ta0.8Hf0.2C and α-SiC phases could be quantiﬁed: the volume fraction of α-SiC was about 27.04 vol%. It was consistent with the calculated value based on Reaction 1. The microstructural evolution of Ta0.8Hf0.2C-27 vol%SiC upon densiﬁcation was revealed in Fig. 3(a–d). The in-situ β-SiC distributed in the matrix based on the XRD pattern at 1650 °C. The morphology of SiC transformed from equiaxed (β-SiC, 1650 °C) to elongated (α-SiC, 2200 °C) grains. This microstructure evolution can be ascribed to the phase transformation and growth behavior . It was reported that the β-α phase transformation of SiC could lead to the in-situ growth of elongated α-SiC grains . As the schematic diagram showed in 3 Journal of the European Ceramic Society xxx (xxxx) xxx–xxx B. Zhang et al. Fig. 3. Backscattered electron images (BSE) of samples sintered at (a)1650 °C and 2200 °C for (b)0 h, (c)1 h and (d) 2 h soaking; (e) schematic diagram showing β-α phase transformation of SiC in Ta0.8Hf0.2C matrix; (f) HRTEM of grain boundary between Ta0.8Hf0.2C and α-SiC grains. smaller crack-opening displacement (U1) than the previous one (U0)  in Fig. 4: adjacent to a crack deﬂection, crack-tip of the branching was observed. When the crack-tip encountered the matrix grains during its propagation course, stress intensity was reduced heavily. Such energy dissipation demonstrated consumed fracture energy and enhanced toughness. Highly-homogenous microstructure and highly-elongated (or platelet-like) α-SiC grains contributed to the higher toughness in the composites: typically, the formation of a tough interfacial interlocking microstructure between α-SiC grains and the matrix led to the improvement of toughness . Besides, residual thermal stress ﬁeld, originated from the thermal expansion mismatch  between SiC and Ta0.8Hf0.2C, might contribute to the crack bridging and deﬂection behaviors. Toughening of our composite was caused by crack-shielding eﬀect, indicating toughness increased as cracks grew, or so-called R-curve behavior . This toughening does little contribute to strength increase. However, a balance between the toughness and strength improvements was achieved by in-situ introduction of α-SiC into the Ta0.8Hf0.2C matrix in our work. The strategy of the strong Ta0.8Hf0.2CSiC bonding and interlocking structure were responsible for the simultaneous improvement of strength and toughness. Table 1 Information of Ta0.8Hf0.2C-27 vol% SiC sintered at 2200 °C for diﬀerent soaking time. Soaking Time (h) 0 1 2 Bulk Density (g/cm3) Relative Density (%) α-SiC Length (μm) Width (μm) Aspect Ratio Matrix Grain Size (μm) 9.71 10.75 11.02 87.2 96.5 99.0 23.6 24.5 26.0 4.0 3.1 1.7 6.0 7.8 15.2 11.2 12.3 13.1 Table 2 Mechanical properties of Ta0.8Hf0.2C based composites sintered at 2200 °C for 2 h. Fraction of SiC HV3/GPa KIC/MPa m1/2 σ/MPa 0vol% 27 vol% 15.9 ± 0.2 16.0 ± 0.1 3.7 ± 0.1 5.4 ± 1.2 359 ± 6 443 ± 22 tortuous and dissipated fracture energy, and therefore toughness was enhanced. Crack branching generated at the strain debonded area of the SiC grains and dissipated fracture energy. The width of propagating cracks became narrower after they passed across SiC particles with Fig. 4. BSE image of indentation crack propagation on polished surface in Ta0.8Hf0.2C-27 vol% SiC sample. 4 Journal of the European Ceramic Society xxx (xxxx) xxx–xxx B. Zhang et al. 4. Conclusions  Ta0.8Hf0.2C-27 vol%SiC composite was pressurless reactive-sintered with a relative density of 99.0% after soaking 2 h at 2200 °C using commercially available TaC, HfSi2 and carbon black. Well-documented cubic phase HfC (PDF#39-1491) and β-SiC (PDF#29-1129, 3C-SiC) peaks were detected after reaction between HfSi2 and carbon at above 1650 °C. Ta0.8Hf0.2C was formed via mutual diﬀusion between raw TaC and in-situ HfC with a lattice parameter of 4.4858 ± 0.0003 Å after sintering at 2200 °C. In-situ β-α SiC phase transformation occurred during sintering. Strain at the β/α interphase was found to inﬂuence the oriented growth of ﬁnal elongated α-SiC grains. The Vickers hardness of Ta0.8Hf0.2C-27 vol%SiC composite was 16.0 ± 0.1 GPa. The ﬂexural strength of the composite increased to 443 ± 22 MPa, which was 23.5% higher than the dense Ta0.8Hf0.2C ceramic. Toughness values of up to 5.4 ± 1.2 MPa m1/2 was measured, which was 45.9% higher than its SiC free counterpart. Formation of a tough interfacial interlocking microstructure between highly-elongated α-SiC grains and the solid solution matrix led to the improvement of toughness.          Acknowledgements  Financial support from National Natural Science Foundation of China (No. 51602325), (No. 51572276), Youth Innovation Promotion Association (CAS), Science Foundation for Youth Scholar of State Key Laboratory of High Performance Ceramics and Superﬁne Microstructures, Shanghai Institute of Ceramics Chinese Academy of Sciences (SKL201602), Scientiﬁc and Technological Innovation Project of Shanghai Institute of Ceramics are gratefully acknowledged.    References    R. Andrievskii, N. Strel’nikova, N. Poltoratskii, E. Kharkhardin, V. Smirnov, Melting point in systems ZrC-HfC, TaC-ZrC, TaC-HfC, Powder Metall. Met. Ceram. 6 (1967) 65–67.  X.-G. Wang, J.-X. Liu, Y.-M. Kan, G.-J. Zhang, Eﬀect of solid solution formation on densiﬁcation of hot-pressed ZrC ceramics with MC (M=V, Nb, and Ta) additions, J. Eur. Ceram. Soc. 32 (2012) 1795–1802.  O. Cedillos-Barraza, S. Grasso, N. Al Nasiri, D.D. Jayaseelan, M.J. Reece, W.E. Lee, Sintering behaviour, solid solution formation and characterisation of TaC, HfC and TaC–HfC fabricated by spark plasma sintering, J. Eur. Ceram. Soc. 36 (2016) 1539–1548.  Q.-J. Hong, A. van de Walle, Prediction of the material with highest known melting point from ab initio molecular dynamics calculations, Phys. Rev. B 92 (2015) 020104.  C. Zhang, A. Gupta, S. Seal, B. Boesl, A. Agarwal, Solid solution synthesis of     5 tantalum carbide-hafnium carbide by spark plasma sintering, J. Am. Ceram. Soc. 100 (2017) 1853–1862. B.C. Schulz, B. Wang, R.A. Morris, D. Butts, G.B. Thompson, Inﬂuence of hafnium carbide on vacuum plasma spray processed tantalum carbide microstructures, J. Eur. Ceram. Soc. 33 (2013) 1219–1224. B. Zhang, J. Yin, J. Chen, Y. Huang, X. Liu, Z. Huang, Pressureless densiﬁcation, microstructure tailoring and properties of Ta0.8Hf0.2C-based composites, J. Eur. Ceram. Soc. 38 (2018) 1227–1236. T. Lu, N. Fleck, The thermal shock resistance of solids, Acta Mater. 46 (1998) 4755–4768. V. Zamora, E. Sánchez‐González, A.L. Ortiz, P. Miranda, F. Guiberteau, Hertzian indentation of a ZrB2–30% SiC ultra‐high‐temperature ceramic up to 800°C in air, J. Am. Ceram. Soc. 93 (2010) 1848–1851. P.I. Pelissari, F. Bouville, V.C. Pandolfelli, D. Carnelli, F. Giuliani, A.P. Luz, E. Saiz, A.R. Studart, Nacre-like ceramic refractories for high temperature applications, J. Eur. Ceram. Soc. 38 (2017) 2186–2193. L. Feng, S.-H. Lee, J. Yin, W. Fahrenholtz, Low-temperature sintering of HfC/SiC nanocomposites using HfSi2-C additives, J. Am. Ceram. Soc. 99 (8) (2016) 2632–2638. A.G. EVans, E.A. Charles, Fracture toughness determinations by indentation, J. Am. Ceram. Soc. 59 (7-8) (1976) 371–372. E. Simonenko, N. Ignatov, N. Simonenko, Y.S. Ezhov, V. Sevastyanov, N. Kuznetsov, Synthesis of highly dispersed super-refractory tantalum-zirconium carbide Ta4ZrC5 and tantalum-hafnium carbide Ta4HfC5 via sol-gel technology, Russ. J. Inorg. Chem. 56 (2011) 1681–1687. J. Yin, H. Zhang, Y. Yan, Z. Huang, X. Liu, D. Jiang, High toughness in pressureless densiﬁed ZrB2-based composites co-doped with boron–titanium carbides, Scripta Mater. 66 (2012) 523–526. E. Rudy, H. Nowotny, Untersuchungen im system hafnium-tantal-kohlenstoﬀ, Monatsh. Chem. 94 (3) (1963) 507–517. G. Arslan, F. Kara, S. Turan, Quantitative X-ray diﬀraction analysis of reactive inﬁltrated boron carbide–aluminium composites, J. Eur. Ceram. Soc. 23 (2003) 1243–1255. H.-G. An, Y.-W. Kim, J.-G. Lee, Eﬀect of initial α-phase content of SiC on microstructure and mechanical properties of SiC–TiC composites, J. Eur. Ceram. Soc. 21 (2001) 93–98. N.P. Padture, In situ-toughened silicon carbide, J. Am. Ceram. Soc. 77 (1994) 519–523. L. Ogbuji, T. Mitchell, A. Heuer, S. Shinozaki, The β→α transformation in polycrystalline SiC: IV, a comparison of conventionally sintered, hot‐pressed, reaction‐sintered, and chemically vapor‐deposited samples, J. Am. Ceram. Soc. 64 (1981) 100–105. N.P. Padture, B.R. Lawn, Toughness properties of a silicon carbide with an in situ induced heterogeneous grain structure, J. Am. Ceram. Soc. 77 (1994) 2518–2522. L. Pienti, D. Sciti, L. Silvestroni, A. Cecere, R. Savino, Ablation tests on HfC-and TaC-based ceramics for aeropropulsive applications, J. Eur. Ceram. Soc. 35 (2015) 1401–1411. L. Xu, C. Huang, H. Liu, B. Zou, H. Zhu, G. Zhao, J. Wang, In situ synthesis of ZrB2–ZrCx ceramic tool materials toughened by elongated ZrB2 grains, Mater. Des. 49 (2013) 226–233. S.C. Zhang, G.E. Hilmas, W.G. Fahrenholtz, Pressureless sintering of ZrB2–SiC ceramics, J. Am. Ceram. Soc. 91 (2008) 26–32. M. Taya, S. Hayashi, A.S. Kobayashi, H. Yoon, Toughening of a particulate‐reinforced ceramic‐matrix composite by thermal residual stress, J. Am. Ceram. Soc. 73 (1990) 1382–1391. T. Ohji, Y.K. Jeong, Y.H. Choa, K. Niihara, Strengthening and toughening mechanisms of ceramic nanocomposites, J. Am. Ceram. Soc. 81 (1998) 1453–1460.