- Email: [email protected]
Fabrication and properties of TiB2-based cermets by spark plasma sintering with CoCrFeNiTiAl high-entropy alloy as sintering aid
Available online at www.sciencedirect.com
ScienceDirect Journal of the European Ceramic Society 35 (2015) 879–886
Fabrication and properties of TiB2-based cermets by spark plasma sintering with CoCrFeNiTiAl high-entropy alloy as sintering aid Wei Ji, Jinyong Zhang, Weimin Wang, Hao Wang, Fan Zhang, Yucheng Wang, Zhengyi Fu ∗ State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, 122 No, Luoshi Road, Hongshan District, Wuhan 430070, Hubei, China Received 19 May 2014; received in revised form 17 October 2014; accepted 22 October 2014 Available online 4 November 2014
Abstract Equiatomic CoCrFeNiTiAl high-entropy alloy (HEA) fabricated by mechanical alloying (MA) was used as sintering aid for the densification of TiB2 by spark plasma sintering (SPS). The wettability between HEA and TiB2 , the phases, microstructures and mechanical properties of the cermets were investigated. In the wetting analysis, excellent properties with a contact angle of 16.73◦ at 1420 ◦ C was confirmed via sessile drop method. For the TiB2 -based ceramics, no brittle phases can be detected in the XRD patterns. TEM observation revealed that there existed a narrow grain boundary between TiB2 and the HEA. A typical structure of HEA can be observed in the liquid phase. An optimized sintering temperature of 1600 ◦ C was determined for the preparation of the TiB2 -5 wt.% HEA ceramics with the combination of mechanical properties, including the relative density of 99.12%, the Vickers hardness of 2356 HV5 and the flexural strength of 800.2 MPa. © 2014 Elsevier Ltd. All rights reserved. Keywords: Titanium diboride; High-entropy alloy; Spark plasma sintering; Wettability; Microstructure
1. Introduction Titanium diboride (TiB2 ) has attracted considerable attentions because of its attractive combination of properties, including high melting point (3225 ◦ C), hardness (30 GPa), elastic modules (more than 500 GPa), good thermal and electrical conductivity and excellent chemical stability. These properties make it a fascinating material for various applications, such as high temperature structural components, cutting tools, armors, wear-resistant parts, conductive coatings and aluminum evaporation boats, etc.1 However, the applications of monolithic TiB2 are limited due to the difficulties in obtaining highly dense bulk materials even by pressure-assisted sintering method. Its strong covalent bonding, low self-diffusion coefficient and oxygenrich-layer (mainly TiO2 and B2 O3 ) covered surface are the main reasons that make it difficult to be sintered.2,3 To overcome these difficulties, different additives have been used to improve the
∗
Corresponding author. Tel.: +86 027 87865484; fax: +86 027 87215421. E-mail address: [email protected] (Z. Fu).
http://dx.doi.org/10.1016/j.jeurceramsoc.2014.10.024 0955-2219/© 2014 Elsevier Ltd. All rights reserved.
sinterability of TiB2 . Generally, the available additives can be classified into two groups: nonmetallic and metallic additives. For nonmetallic additives, TiN, AlN, Si3 N4 , MoSi2 , TiSi2 , etc., have been used,4–7 in which high sintering temperature is still needed. For metallic additives, transition metals such as iron, nickel, and titanium have been used.8–11 The good wettability of these metals with TiB2 remarkably decreases the sintering temperature.12 However, the formation of secondary borides (M2 B, M23 B6 ) embrittles these products. Based on the above considerations, it is important to develop new sintering additive for TiB2 . High-entropy alloy (HEA) is one kind of the new sintering additives for ceramics,13 which could lower the sintering temperature and inhibit the grain growth of ceramic, and possess excellent wettability with ceramic. A HEA is originally defined as an alloy system composed of at least five principal elements in an equimolar or near equimolar ratio, with a small difference in atom radii (<15%) and concentration of each element varying from 5 to 35 at.%. The high mixing entropy of multiprinciple elements will induce lattice distortion and sluggish cooperative diffusion. As a consequence, HEA often plays as
880
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
simple solid-solutions and amorphous structure rather than intermetallics. With proper composition designing, the HEA exhibits high hardness, excellent ductility as well as promising resistances to wear, oxidation and corrosion.14,15 The HEA powder has been successfully prepared by mechanical alloying (MA) with fine particle size and good chemical uniformity, so it is suitable to be chosen as a sintering additive for ceramics.16 However, so far rare work relating to the densification of ceramic with HEAs as sintering aids has been done. So it is of great meaning to expand this research scope. The target of this work is to fabricate dense TiB2 -based ceramics using HEA as sintering additive. Based on this aim, an equiatomic CoCrFeNiTiAl HEA was fabricated by MA and the wettability between as-milled HEA and TiB2 was first investigated. Then the TiB2 –HEA ceramics were prepared via spark plasma sintering (SPS) method. Additionally, the microstructures and mechanical properties of the bulks were studied. 2. Experiment details
cut from a dense monolithic TiB2 bulk and sectioned into a wafer with a diameter of 20 mm. The HEA powders were consolidated into the shape of cylinder and positioned in the center of the TiB2 substrate. The test proceeded at a temperature of 1420 ◦ C, which is higher than the melting point of the alloy.15 2.3. Phase composition and microstructure The crystal structure of the as-preserved specimen was characterized by X-ray diffractometer (XRD, Rigaku Ultima III) with Cu K␣ radiation. The fracture surface of the composites was observed using scanning electron microscopy (SEM, Hitachi 3400). A thin foil of sintered material was prepared by mechanical thinning followed by ion milling and observed using a transmission electron microscopy (TEM, JEOL JEM-2010HT) equipped with energy-dispersive X-ray spectroscopy (EDS, EDAX) and selected area electron diffraction (SAED). The more detailed information of the microstructure was observed using a high resolution transmission electron microscopy (HRTEM, JEOL JEM-2010FEF).
2.1. Sample preparation 2.4. Mechanical properties Commercially available cobalt (99% pure), chromium (99.9% pure), iron (98% pure), nickel (99.5% pure), titanium (99.9% pure), aluminum (99.5% pure) and titanium diboride (98% pure, −3 m, Hubei Duobo New Ceramic Materials Co., Ltd., Wuhan, China) powders were used as starting materials. The major impurities of as synthesized TiB2 include oxygen, carbon and nitrogen. All the metal powder were supplied by Sinopharm Chemical Reagent Co., Ltd., Shanghai, China and granulated with a 325 mesh sieve. The elemental powders of Co, Cr, Fe, Ni, Ti and Al were mixed in equiatomic composition and milled in a planetary ball-miller (QM-BP, Nanjing NanDa Instrument Plant, Nanjing, China) for 60 h at 250 rpm in an argon atmosphere. High performance stainless steel vials and balls were utilized as the milling media with a ball-to-powder mass ratio of 15:1. N-heptane acts as the processing controlling agent (PCA) to avoid cold welding as well as to prevent the alloy from oxidizing. Powder mixtures of 95 wt.% TiB2 and 5 wt.% CoCrFeNiTiAl HEA were obtained by a wet ball-miller (GMJ/B, Xianyang JinHong general machinery Co., Ltd., Xianyang, China) in polyethylene jars with agate balls and ethyl alcohol as media for 24 h. The slurry was dried in a rotary evaporator at 70 ◦ C and granulated with a 200 mesh sieve. The TiB2 -based ceramics were obtained by SPS (Dr. Sinter-3.20MK II, SCM) using a graphite die with a diameter of 20 mm, under temperatures between 1300 and 1700 ◦ C with a stepsize of 100 ◦ C. All SPS experiments were carried out with a maximum applied pressure of 30 MPa, heating rate of 100 ◦ C/min and soaking time of 5 min in vacuum. 2.2. Wettability analysis A high-temperature microscopy (EM-201, Leitz) was used to observe the contact angle between CoCrFeNiTiAl HEA and TiB2 substrate by sessile drop method.17 The TiB2 substrate was
Specimens for mechanical testing were cut from the SPS-ed cermets. Each specimen was polished with diamond suspension down to 0.25 m. The edges were chamfered to minimize stress concentration effects. The Vickers hardness was measured by a Vickers hardness tester (Wolpert-430SV) with a load of 5 kg for 15 s. The flexural strength measurements were carried out via three point bending test on 2 mm × 3 mm × 18 mm bars in a ceramic test system (MTS 810, MTS) with a span of 15 mm at a crosshead speed of 0.5 mm/min. 3. Results and discussion The characteristics of as-milled equiatomic CoCrFeNiTiAl high-entropy alloy are shown in Fig. 1. The XRD patterns shown in Fig. 1(a) reveal that a phase of CoCrFeNiTiAl HEA is formed within a milling time of 30 h. The primitive blending powder includes diffraction patterns of all alloying elements. After MA for 6 h, the diffraction peaks of the principle elements can still be observed with a dramatic decrease in intensity. With prolonged milling time, peak broadening is obvious and some peaks become invisible after 18 h of milling. As the milling time increases to 30 h, only the 3 most intensive peaks of a BCC structure ((1 1 0), (2 0 0), (2 1 1)) can be identified, which implies the formation of a simple solid solution. Further extended duration up to 60 h results in no obvious change in the XRD patterns. Throughout the milling process, the decrease in intensity, broadening of the peak and its subsequent disappearance may result from the three following factors: refined crystal size, high lattice strain and decreased crystallinity.18 The nanocrystalline nature of CoCrFeNiTiAl HEA after 60 h MA has been further characterized by TEM, as shown in Fig. 1(b). The crystal size measured in the bright field image is approximately 10 nm. The rings in the SAED pattern indicate that the nanocrystalline HEA powder prepared by milling for 60 h has a crystal structure of BCC,
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
881
Fig. 1. The characteristics of as-milled equiatomic CoCrFeNiTiAl high-entropy alloy: (a) XRD patterns with different milling time from 0 h to 60 h; (b) TEM images and SAED patterns of 60 h milled HEA powder.
which is in agreement with the XRD analysis. This result confirms that the CoCrFeNiTiAl high-entropy alloy with a structure of simple BCC solid solution has been successfully fabricated by mechanical alloying. Fig. 2 shows the image of sessile drop of HEA on the pure TiB2 substrate at 1420 ◦ C. The white peak on the horizontal line was molten HEA and the rest part of white area was TiB2 . It can be distinctly observed that the mean contact angle is equal to 16.73◦ and the profile of droplets is stable within experimental duration. Compared with the available literature,19 the result indicates good wettability of TiB2 ceramic by molten CoCrFeNiTiAl HEA, which could promote the sinterability of TiB2 ceramics. Wetting property between metals and ceramics is one of the most important aspects in the fabrication of metal-ceramic composites by liquid sintering routing. However, rare research has been reported on the wetting property between ceramic and HEA. So this groundbreaking work is irradiative for further investigation in this field. The XRD analysis of mixed powder and SPS-ed TiB2 ceramics with 5 wt.% CoCrFeNiTiAl HEA as additive were carried out
and the results are presented in Fig. 3. Only a TiB2 phase was detected in the mixed powder, while the predominant percentage of TiB2 and an inconspicuous amount of Ti(C, N, O) and Al2 O3 phases can be detected in the TiB2 -based ceramics. The formation of Ti(C, N, O) and Al2 O3 results from high-temperature reactions between the impurities on the surface of raw TiB2 powders and the elements exsoluted from the HEA. These reactions can be considered as a routing to collect impurities in raw materials, which promote the sinterability of the ceramic. According to the patterns, no obvious peaks corresponding to the HEA can be observed. The invisibility of HEA peaks can be ascribed to: (1) the main peaks of as-milled CoCrFeNiTiAl HEA and TiB2 are almost overlapped at a 2θ of about 44.5◦ ; (2) the small content of HEA addition (5 wt.%) and (3) the low crystallinity of HEA owing to the mechanical alloying routing. For the former two factors, it is very hard to be distinguished for such closed peaks in the XRD patterns, let alone a small fraction of second phase. Moreover, the alloy powder prepared by MA surely
Fig. 2. Image for sessile droplets of molten CoCrFeNiTiAl HEA on TiB2 substrate at 1420 ◦ C.
Fig. 3. The XRD patterns of mixed powder and SPS-ed TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics under different sintering temperature.
882
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
consists of random solid solution, nano-particles even amorphous phase with severe lattice distortion.20 As a consequence, the intensity of X-ray diffraction peaks is not strong enough to be identified. The existence of HEA in the compound will be further confirmed hereinafter. It is worth pointing out that there is no peak belongs to the brittle secondary borides (M2 B, M23 B6 ) in the XRD patterns, which indicates that the products possess promising mechanical properties. Fig. 4 exhibits the SEM images of TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics prepared at different sintering temperatures. The secondary electron SEM images of fracture surface are shown in Fig. 4(a)–(e). Pores can be observed in Fig. 4(a), which is due to the low density of TiB2 under low temperature. With the increase of sintering temperature, the densification of TiB2 was remarkably enhanced, as shown in Fig. 4(b)–(e). However, coarser grains are also observed at the higher temperature. A widely accepted view is that ceramics materials with higher density and smaller grain size have better mechanical properties. So it is expected to achieve an optimized sintering process in this sintering system. From the fracture surface of the samples, it can be seen that fracture occurred mainly by the intergranular mode at 1300 ◦ C as a result of the high porosity. With the increase of temperature, the fracture mode evolves to the combination of transgranular and intergranular. The transgranular fracture mode may result from the strong bonding between TiB2 and HEA. A high magnification backscattered electron SEM image of 1600 ◦ C sintered TiB2 –HEA ceramics was shown in Fig. 4(f). Compared with the secondary electron SEM images, a light phase in liquid structure can be observed in the grain boundaries, which can be presumed as HEA preliminarily. It can be seen that the shape of the liquid phase is dependent on the profile of adjacent TiB2 grains. Owing to the excellent fluidity at high temperature, the liquid phase could play an important role in joining the TiB2 grains and accelerating the densification of ceramics. Moreover, the morphology also indicates a good wettability between TiB2 and HEA. For further investigating the microstructure and the mechanisms of the densification, TEM analysis was performed on the TiB2 –HEA ceramic sintered at 1600 ◦ C. Fig. 5 shows a bright field TEM image with SAED patterns of the grain and EDS patterns of liquid phase. Analysis of the SAED patterns shown in Fig. 5(b) reveals that the incident beam is paralleled to the [1¯ ¯1 2 2] zone axis of the TiB2 . The EDS patterns shown in Fig. 5(c) reveal that the liquid phase only contains the elements of Co, Cr, Fe, Ni, Ti and Al, and no B element can be detected, which evidences that the liquid phase is HEA. In addition, only a small fraction of Al can be observed, which may be related to the reactions with the oxygen-rich-layer of raw TiB2 . The HRTEM images of the boundary between ceramic grains and liquid phase are shown in Fig. 6. The selected region is magnified in Fig. 6(b) with distinct observation of a narrow grain boundary between a TiB2 ceramic grain (d = 0.204 nm, [1 0 1¯ 1] zone axis) and HEA liquid phase. This kind of tight and narrow grain boundary structure with low strain energy indicates excellent physical compatibility between TiB2 and HEA. It is well known that the grain boundary between matrix and second phases has a significant effect on the properties of the
composites, especially the mechanical properties. Therefore it is predictable that excellent mechanical properties can be obtained for the cermets. Moreover, some dispersed nanoparticles are observed to be embedded in the amorphous matrix, which is similar with the typical morphology of HEAs.21 The existence of amorphous phase can be explained by two points. First, amorphous phase would be produced during mechanical alloying. As we known, the MA processing is in favor of amorphization. Second, according to Ref. [14], there are four core effects for HEAs, which are: (1) high entropy effect; (2) sluggish diffusion effect; (3) severe lattice distortion effect and (4) cocktail effects. The effects of (1), (2), and (3) can slow down the atom diffusion rate and inhibit the crystallization. So for a HEA system with high entropy of 1.79R and severe lattice distortion, some of the amorphous phase could exist in the HEA region of the cermets. The microstructures of certain HEAs are usually very complicated, which can include nano-precipitates, ordered or disordered solid-solution phases and amorphous phases. Precipitation of nanosized particles is a novel feature of HEAs because of the effect of high mixing entropy. From the viewpoint of kinetics, long-range diffusion for phase separation is sluggish in the HEAs with multi-principal elements. Difficulties in substitutional diffusion of elements in these alloys and interactions among interdiffusing species during partitioning lower the rates of nucleation and growth, and induce the formation of nanosized crystallites. It is of great meaning and interest to investigate the coexistent of amorphous phase and nanoparticles in the liquid phase. To confirm the phase composition, the sample was further ion milled and a more identifiable HRTEM image with corresponding selected area fast Fourier transform (FFT) image of the liquid phase are shown in Fig. 7. It clearly shows that nanoparticles with size of 10 nm are embedded in the amorphous matrix and these nanoparticles are identified as a simple FCC structure by the FFT image.12 The particle size is in good agreement with that of the mechanical alloyed CoCrFeNiTiAl HEA. However, the crystal structure is different from the mechanical alloyed CoCrFeNiTiAl HEA. According to available literatures,16,22 the BCC structured HEA synthesized by mechanical alloying is a metastable state because of the non-equilibrium processing. This metastable structure will convert to a more stable phase (FCC structure here) at high temperature of annealing. That is why the HEA in the cermet is in FCC structure. In addition, FCC structure has higher atomic packing efficiency than BCC structure (FCC 74% and BCC 68%). According to Refs. [23,24], the dense packing with a large difference in the atomic size would make the alloy melt with high viscosity and low atomic mobility, which greatly reduce the nucleation rate and growth velocity, and facilitate the amorphization, even at low cooling rate. Some defects as twining crystal and lattice distortion can be observed in the nanoparticles, as shown in Fig. 7 (b)–(c), which are caused by the high energy of mechanical alloying and the large atomic radius difference between so many elements.25 This kind of structure is in favor of promoting the strength of the grain boundary phase in the cermets. The effect of the sintering temperature on mechanical properties of the TiB2 –HEA ceramics prepared by SPS is shown in
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
883
Fig. 4. The SEM images of TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics: ((a)–(e)) secondary electron images at different temperatures: (a)1300 ◦ C; (b)1400 ◦ C; (c) 1500 ◦ C; (d) 1600 ◦ C; (e) 1700 ◦ C; (f) back-scattered electron image at 1600 ◦ C.
Fig. 5. TEM images of the TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics sintered at 1600 ◦ C: (a) bright field TEM image with (b) corresponding SAED and (c) EDS patterns.
884
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
Fig. 6. The HRTEM images of the interface between ceramic grains and liquid phase for the TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics sintered at 1600 ◦ C.
Fig. 8. It reveals that the sintering temperature has a remarkable influence on the relative density, Vickers hardness and flexural strength of TiB2 –HEA ceramics. The relative density increases with increasing temperature from 1400 to 1600 ◦ C. However, further increase of sintering temperature has no positive influence on the densification of the samples. The similar trends were observed in the Vickers hardness and flexural strength, as seen in Fig. 8. The increases in hardness and strength should be mainly attributed to the improvement in bulk densities and the enhancement in bond strength between TiB2 and HEA. Grain growth restricts the further improvement of the mechanical properties at the higher temperature. The TiB2 –HEA ceramics sintered at 1600 ◦ C exhibits the attractive combination of mechanical properties, including a relative density of 99.12%,
a Vickers hardness of 2356 HV5 and a flexural strength of 800.2 MPa. A Comparison of the properties of the TiB2 -5 wt.% HEA cermets with other TiB2 -based ceramics is shown in Table 1. For monolithic TiB2 ceramics, a SPS temperature of higher than 1800 ◦ C is required for attaining high relative density (>97%).10 However, it is demonstrated by our work that the highly-dense ceramics can also be fabricated at a decreased temperature (1600 ◦ C). Besides, the TiB2 -based ceramics with a sintering aid of CoCrFeNiTiAl HEA processes a better combination of dense microstructure and excellent mechanical properties. The improvement in microstructures and mechanical properties of the ceramics in the present work due to three points: First, for most of metallic sintering aid, the formation of secondary
Fig. 7. The HRTEM images of the liquid phase for the TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics sintered at 1600 ◦ C: (a) bright field TEM image with corresponding selected area fast Fourier transform image; ((b)–(d)) magnified nano-particles in the liquid phase.
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
885
Table 1 Comparison of the properties of the TiB2 -5 wt.% HEA cermets with other TiB2 -based ceramics. Material (wt.%) Monolithic TiB2 TiB2 –5 B4 C–0.5 Fe TiB2 –2.5 Ti TiB2 –10 MoSi2 TiB2 –5 Si3 N4 TiB2 –5 AlN TiB2 –5 HEA
Processing 1800 ◦ C,
HP, 1 h, 30 MPa PS, 2000 ◦ C, 30 min SPS, 1650 ◦ C, 5 min, 50 MPa HP, 1700 ◦ C, 1 h, 30 MPa HP, 1800 ◦ C, 1 h, 30 MPa HP, 1800 ◦ C, 1 h, 30 MPa SPS, 1600 ◦ C, 5 min, 30 MPa
Relatively density (%)
Flexural strength (MPa)
Vickers hardness (GPa)
Reference
96.1 95 99.1 96.7 97.5 98 99.12
367.1 (4-P) 400 (3-P) 558 (3-P) 267.8 (3-P) 500 (3-P) 650 (4-P) 800.2 (3-P)
24 (98 N) – 26.8 (50 N) 22.97 (49 N) 22 (49 N) 21.5 (49 N) 23.56 (49 N)
[6] [11] [10] [6] [5] [4] Present work
microstructure and excellent mechanical properties, including a relative density of 99.12%, a Vickers hardness of 2356 HV5 and a flexural strength of 800.02 MPa. Acknowledgments This work was financially supported by the National Natural Science Foundation of China and the international cooperation project from Ministry of Science and Technology of China. References
Fig. 8. Mechanical properties of TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics at different sintering temperature.
borides (M2 B or M23 B6 ) embrittles the products. But in our work, no brittle phases were detected in the cermets. Second, for the crystal structure of HEA, the FCC-structured HEAs have high plasticity and relatively low yield strength, while the BCC-structured HEAs usually have very high yield strengths and limited plasticity.26 These characteristics reveal that one kind of the toughening mechanisms for this ceramic is ductile phase toughening. Third, SEM images (shown in Fig. 4) illustrate that the fracture mode changes from intergranular fracture to a combination of transgranular and intergranular fracture with the increase of relatively density. The morphologies reveal good wetting ability and bonding between the phases and also indicates crack deflection of the dense ceramic, which also promote the mechanical properties. 4. Conclusion Based on the above results and discussion, we can draw a conclusion as followed: (a) Excellent wettability (a contact angle of 16.73◦ ) between TiB2 ceramic and CoCrFeNiTiAl high-entropy alloy was detected by a high-temperature microscopy. (b) TEM observations reveal tight and narrow grain boundary between TiB2 ceramic and HEA liquid phase, which indicates excellent physical compatibility between the two components. (c) The TiB2 -based ceramic sintered at a low temperature of 1600 ◦ C provides the attractive combination of dense
1. Wen G, Li S, Zhang B, Guo Z. Reaction synthesis of TiB2 –TiC composites with enhanced toughness. Acta Mater 2001;49:1463–70. 2. Raju G, Mukhopadhyay A, Biswas K, Basu B. Densification and hightemperature mechanical properties of hot pressed TiB2 -(0–10 wt.%) MoSi2 composites. Scr Mater 2009;61:674–7. 3. Wang W, Fu Z, Wang H, Yuan R. Influence of hot pressing sintering temperature and time on microstructure and mechanical properties of TiB2 ceramics. J Eur Ceram Soc 2002;22:1045–9. 4. Li L, Kim H, Kang E. Sintering and mechanical properties of titanium diboride with aluminum nitride as a sintering aid. J Eur Ceram Soc 2002;22:973–7. 5. Park J, Koh Y, Kim H, Hwang C, Kang E. Densification and mechanical properties of titanium diboride with silicon nitride as a sintering aid. J Am Ceram Soc 1999;82:3037–42. 6. Mukhopadhyay A, Raju G, Basu B, Suri A. Correlation between phase evolution, mechanical properties and instrumented indentation response of TiB2 -based ceramics. J Eur Ceram Soc 2009;29:505–16. 7. Zheng L, Li F, Zhou Y. Preparation, microstructure, and mechanical properties of TiB2 using Ti3 AlC2 as a sintering aid. J Am Ceram Soc 2012;95:2028–34. 8. Barandika M, Sanchez J, Rojo T, Cortes R, Castro F. Fe-Ni-Ti binder phases for TiB2 -based cermets: a thermodynamic approach. Scr Mater 1998;39:1395–400. 9. Einarsrud M, Hagen E, Pettersen G, Grande T. Pressureless sintering of titanium diboride with nickel, nickel boride and iron additives. J Am Ceram Soc 1997;80:3013–20. 10. Zhang Z, Shen X, Wang F, Lee S, Fan Q, Cao M. Low-temperature densification of TiB2 ceramic by the spark plasma sintering process with Ti as a sintering aid. Scr Mater 2012;66:167–70. 11. Kang E, Jang C, Lee C, Kim C, Kim D. Effect of iron and boron carbide on the densification and mechanical properties of titanium diboride ceramics. J Am Ceram Soc 1989;72:1868–72. 12. Zhang J, Fu Z, Wang W, Wang H, Min X. Wettability between TiB2 ceramic and metals. Acta Metall Sin 1999;12:395–400. 13. Zhu G, Liu Y, Ye J. Fabrication and properties of Ti(C, N)-based cermets with multi-component AlCoCrFeNi high-entropy alloys binder. Mater Lett 2013;113:80–2. 14. Yeh J, Chen S, Lin S, Gan J, Chin T, Shun T, et al. Nanostructured highentropy alloys with multiple principal elements: novel alloy design concepts and outcomes. Adv Eng Mater 2004;6:299–303.
886
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
15. Otto F, Yang Y, Bei H, George E. Relative effects of enthalpy and entropy on the phase stability of equiatomic high-entropy alloys. Acta Mater 2013;61:2628–38. 16. Ji W, Fu Z, Wang W, Wang H, Zhang J, Wang Y, et al. Mechanical alloying synthesis and spark plasma sintering consolidation of CoCrFeNiAl highentropy alloy. J Alloy Compd 2014;589:61–6. 17. Zheng X, Shen P, Han X, Lin Q, Qiu F, Zhang Y, et al. Wettability and reactivity between B4 C and Molten Zr55 Cu30 Al10 Ni5 metallic glass alloy. Mater Chem Phys 2009;117:377–83. 18. Yeh J, Chang S, Hong Y, Chen S, Lin S. Anomalous decrease in X-ray diffraction intensities of Cu–Ni–Al–Co–Cr–Fe–Si alloy systems with multiprincipal elements. Mater Chem Phys 2007;103:41–6. 19. Aizenshtein M, Froumin N, Frage N. The nature of TiB2 wetting by Cu and Au. J Mater Eng Perform 2012;21:655–9. 20. Varalakshmi S, Kamaraj M, Murty B. Synthesis and characterization of nanocrystalline AlFeTiCrZnCu high entropy solid solution by mechanical alloying. J Alloy Compd 2008;460:253–7.
21. Zhang X, Zhang Y, Qiao Y, Chen G. Novel microstructure and properties of multicomponent CoCrCuFeNiTix alloys. Intermetallics 2007;15: 357–62. 22. Ji W, Wang W, Wang H, Zhang J, Wang Y, Zhang F, et al. Alloying behavior and novel properties of CoCrFeNiMn high-entropy alloy fabricated by mechanical alloying and spark plasma sintering. Intermetallics 2015;56:24–7. 23. Li Y, Guo Q, Kalb, Thompson C. Matching glass-forming ability with the density of the amorphous phase. Science 2008;322:1816–9. 24. Li R, Pang S, Men H, Ma C, Zhang T. Formation and mechanical properties of (Ce–La–Pr–Nd)–Co–Al bulk glassy alloys with superior glass-forming ability. Scr Mater 2006;54:1123–6. 25. Fu Z, Chen W, Fang S, Zhang D, Xiao H, Zhu D. Alloying behavior and deformation twinning in a CoNiFeCrAl0.6 Ti0.4 high entropy alloy processed by spark plasma sintering. J Alloy Compd 2013;553:316–23. 26. Zhang Y, Zuo T, Tang Z, Gao M, Dahmen K, Liaw P, et al. Microstructures and properties of high-entropy alloys. Prog Mater Sci 2014;61:1–93.
ScienceDirect Journal of the European Ceramic Society 35 (2015) 879–886
Fabrication and properties of TiB2-based cermets by spark plasma sintering with CoCrFeNiTiAl high-entropy alloy as sintering aid Wei Ji, Jinyong Zhang, Weimin Wang, Hao Wang, Fan Zhang, Yucheng Wang, Zhengyi Fu ∗ State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, 122 No, Luoshi Road, Hongshan District, Wuhan 430070, Hubei, China Received 19 May 2014; received in revised form 17 October 2014; accepted 22 October 2014 Available online 4 November 2014
Abstract Equiatomic CoCrFeNiTiAl high-entropy alloy (HEA) fabricated by mechanical alloying (MA) was used as sintering aid for the densification of TiB2 by spark plasma sintering (SPS). The wettability between HEA and TiB2 , the phases, microstructures and mechanical properties of the cermets were investigated. In the wetting analysis, excellent properties with a contact angle of 16.73◦ at 1420 ◦ C was confirmed via sessile drop method. For the TiB2 -based ceramics, no brittle phases can be detected in the XRD patterns. TEM observation revealed that there existed a narrow grain boundary between TiB2 and the HEA. A typical structure of HEA can be observed in the liquid phase. An optimized sintering temperature of 1600 ◦ C was determined for the preparation of the TiB2 -5 wt.% HEA ceramics with the combination of mechanical properties, including the relative density of 99.12%, the Vickers hardness of 2356 HV5 and the flexural strength of 800.2 MPa. © 2014 Elsevier Ltd. All rights reserved. Keywords: Titanium diboride; High-entropy alloy; Spark plasma sintering; Wettability; Microstructure
1. Introduction Titanium diboride (TiB2 ) has attracted considerable attentions because of its attractive combination of properties, including high melting point (3225 ◦ C), hardness (30 GPa), elastic modules (more than 500 GPa), good thermal and electrical conductivity and excellent chemical stability. These properties make it a fascinating material for various applications, such as high temperature structural components, cutting tools, armors, wear-resistant parts, conductive coatings and aluminum evaporation boats, etc.1 However, the applications of monolithic TiB2 are limited due to the difficulties in obtaining highly dense bulk materials even by pressure-assisted sintering method. Its strong covalent bonding, low self-diffusion coefficient and oxygenrich-layer (mainly TiO2 and B2 O3 ) covered surface are the main reasons that make it difficult to be sintered.2,3 To overcome these difficulties, different additives have been used to improve the
∗
Corresponding author. Tel.: +86 027 87865484; fax: +86 027 87215421. E-mail address: [email protected] (Z. Fu).
http://dx.doi.org/10.1016/j.jeurceramsoc.2014.10.024 0955-2219/© 2014 Elsevier Ltd. All rights reserved.
sinterability of TiB2 . Generally, the available additives can be classified into two groups: nonmetallic and metallic additives. For nonmetallic additives, TiN, AlN, Si3 N4 , MoSi2 , TiSi2 , etc., have been used,4–7 in which high sintering temperature is still needed. For metallic additives, transition metals such as iron, nickel, and titanium have been used.8–11 The good wettability of these metals with TiB2 remarkably decreases the sintering temperature.12 However, the formation of secondary borides (M2 B, M23 B6 ) embrittles these products. Based on the above considerations, it is important to develop new sintering additive for TiB2 . High-entropy alloy (HEA) is one kind of the new sintering additives for ceramics,13 which could lower the sintering temperature and inhibit the grain growth of ceramic, and possess excellent wettability with ceramic. A HEA is originally defined as an alloy system composed of at least five principal elements in an equimolar or near equimolar ratio, with a small difference in atom radii (<15%) and concentration of each element varying from 5 to 35 at.%. The high mixing entropy of multiprinciple elements will induce lattice distortion and sluggish cooperative diffusion. As a consequence, HEA often plays as
880
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
simple solid-solutions and amorphous structure rather than intermetallics. With proper composition designing, the HEA exhibits high hardness, excellent ductility as well as promising resistances to wear, oxidation and corrosion.14,15 The HEA powder has been successfully prepared by mechanical alloying (MA) with fine particle size and good chemical uniformity, so it is suitable to be chosen as a sintering additive for ceramics.16 However, so far rare work relating to the densification of ceramic with HEAs as sintering aids has been done. So it is of great meaning to expand this research scope. The target of this work is to fabricate dense TiB2 -based ceramics using HEA as sintering additive. Based on this aim, an equiatomic CoCrFeNiTiAl HEA was fabricated by MA and the wettability between as-milled HEA and TiB2 was first investigated. Then the TiB2 –HEA ceramics were prepared via spark plasma sintering (SPS) method. Additionally, the microstructures and mechanical properties of the bulks were studied. 2. Experiment details
cut from a dense monolithic TiB2 bulk and sectioned into a wafer with a diameter of 20 mm. The HEA powders were consolidated into the shape of cylinder and positioned in the center of the TiB2 substrate. The test proceeded at a temperature of 1420 ◦ C, which is higher than the melting point of the alloy.15 2.3. Phase composition and microstructure The crystal structure of the as-preserved specimen was characterized by X-ray diffractometer (XRD, Rigaku Ultima III) with Cu K␣ radiation. The fracture surface of the composites was observed using scanning electron microscopy (SEM, Hitachi 3400). A thin foil of sintered material was prepared by mechanical thinning followed by ion milling and observed using a transmission electron microscopy (TEM, JEOL JEM-2010HT) equipped with energy-dispersive X-ray spectroscopy (EDS, EDAX) and selected area electron diffraction (SAED). The more detailed information of the microstructure was observed using a high resolution transmission electron microscopy (HRTEM, JEOL JEM-2010FEF).
2.1. Sample preparation 2.4. Mechanical properties Commercially available cobalt (99% pure), chromium (99.9% pure), iron (98% pure), nickel (99.5% pure), titanium (99.9% pure), aluminum (99.5% pure) and titanium diboride (98% pure, −3 m, Hubei Duobo New Ceramic Materials Co., Ltd., Wuhan, China) powders were used as starting materials. The major impurities of as synthesized TiB2 include oxygen, carbon and nitrogen. All the metal powder were supplied by Sinopharm Chemical Reagent Co., Ltd., Shanghai, China and granulated with a 325 mesh sieve. The elemental powders of Co, Cr, Fe, Ni, Ti and Al were mixed in equiatomic composition and milled in a planetary ball-miller (QM-BP, Nanjing NanDa Instrument Plant, Nanjing, China) for 60 h at 250 rpm in an argon atmosphere. High performance stainless steel vials and balls were utilized as the milling media with a ball-to-powder mass ratio of 15:1. N-heptane acts as the processing controlling agent (PCA) to avoid cold welding as well as to prevent the alloy from oxidizing. Powder mixtures of 95 wt.% TiB2 and 5 wt.% CoCrFeNiTiAl HEA were obtained by a wet ball-miller (GMJ/B, Xianyang JinHong general machinery Co., Ltd., Xianyang, China) in polyethylene jars with agate balls and ethyl alcohol as media for 24 h. The slurry was dried in a rotary evaporator at 70 ◦ C and granulated with a 200 mesh sieve. The TiB2 -based ceramics were obtained by SPS (Dr. Sinter-3.20MK II, SCM) using a graphite die with a diameter of 20 mm, under temperatures between 1300 and 1700 ◦ C with a stepsize of 100 ◦ C. All SPS experiments were carried out with a maximum applied pressure of 30 MPa, heating rate of 100 ◦ C/min and soaking time of 5 min in vacuum. 2.2. Wettability analysis A high-temperature microscopy (EM-201, Leitz) was used to observe the contact angle between CoCrFeNiTiAl HEA and TiB2 substrate by sessile drop method.17 The TiB2 substrate was
Specimens for mechanical testing were cut from the SPS-ed cermets. Each specimen was polished with diamond suspension down to 0.25 m. The edges were chamfered to minimize stress concentration effects. The Vickers hardness was measured by a Vickers hardness tester (Wolpert-430SV) with a load of 5 kg for 15 s. The flexural strength measurements were carried out via three point bending test on 2 mm × 3 mm × 18 mm bars in a ceramic test system (MTS 810, MTS) with a span of 15 mm at a crosshead speed of 0.5 mm/min. 3. Results and discussion The characteristics of as-milled equiatomic CoCrFeNiTiAl high-entropy alloy are shown in Fig. 1. The XRD patterns shown in Fig. 1(a) reveal that a phase of CoCrFeNiTiAl HEA is formed within a milling time of 30 h. The primitive blending powder includes diffraction patterns of all alloying elements. After MA for 6 h, the diffraction peaks of the principle elements can still be observed with a dramatic decrease in intensity. With prolonged milling time, peak broadening is obvious and some peaks become invisible after 18 h of milling. As the milling time increases to 30 h, only the 3 most intensive peaks of a BCC structure ((1 1 0), (2 0 0), (2 1 1)) can be identified, which implies the formation of a simple solid solution. Further extended duration up to 60 h results in no obvious change in the XRD patterns. Throughout the milling process, the decrease in intensity, broadening of the peak and its subsequent disappearance may result from the three following factors: refined crystal size, high lattice strain and decreased crystallinity.18 The nanocrystalline nature of CoCrFeNiTiAl HEA after 60 h MA has been further characterized by TEM, as shown in Fig. 1(b). The crystal size measured in the bright field image is approximately 10 nm. The rings in the SAED pattern indicate that the nanocrystalline HEA powder prepared by milling for 60 h has a crystal structure of BCC,
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
881
Fig. 1. The characteristics of as-milled equiatomic CoCrFeNiTiAl high-entropy alloy: (a) XRD patterns with different milling time from 0 h to 60 h; (b) TEM images and SAED patterns of 60 h milled HEA powder.
which is in agreement with the XRD analysis. This result confirms that the CoCrFeNiTiAl high-entropy alloy with a structure of simple BCC solid solution has been successfully fabricated by mechanical alloying. Fig. 2 shows the image of sessile drop of HEA on the pure TiB2 substrate at 1420 ◦ C. The white peak on the horizontal line was molten HEA and the rest part of white area was TiB2 . It can be distinctly observed that the mean contact angle is equal to 16.73◦ and the profile of droplets is stable within experimental duration. Compared with the available literature,19 the result indicates good wettability of TiB2 ceramic by molten CoCrFeNiTiAl HEA, which could promote the sinterability of TiB2 ceramics. Wetting property between metals and ceramics is one of the most important aspects in the fabrication of metal-ceramic composites by liquid sintering routing. However, rare research has been reported on the wetting property between ceramic and HEA. So this groundbreaking work is irradiative for further investigation in this field. The XRD analysis of mixed powder and SPS-ed TiB2 ceramics with 5 wt.% CoCrFeNiTiAl HEA as additive were carried out
and the results are presented in Fig. 3. Only a TiB2 phase was detected in the mixed powder, while the predominant percentage of TiB2 and an inconspicuous amount of Ti(C, N, O) and Al2 O3 phases can be detected in the TiB2 -based ceramics. The formation of Ti(C, N, O) and Al2 O3 results from high-temperature reactions between the impurities on the surface of raw TiB2 powders and the elements exsoluted from the HEA. These reactions can be considered as a routing to collect impurities in raw materials, which promote the sinterability of the ceramic. According to the patterns, no obvious peaks corresponding to the HEA can be observed. The invisibility of HEA peaks can be ascribed to: (1) the main peaks of as-milled CoCrFeNiTiAl HEA and TiB2 are almost overlapped at a 2θ of about 44.5◦ ; (2) the small content of HEA addition (5 wt.%) and (3) the low crystallinity of HEA owing to the mechanical alloying routing. For the former two factors, it is very hard to be distinguished for such closed peaks in the XRD patterns, let alone a small fraction of second phase. Moreover, the alloy powder prepared by MA surely
Fig. 2. Image for sessile droplets of molten CoCrFeNiTiAl HEA on TiB2 substrate at 1420 ◦ C.
Fig. 3. The XRD patterns of mixed powder and SPS-ed TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics under different sintering temperature.
882
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
consists of random solid solution, nano-particles even amorphous phase with severe lattice distortion.20 As a consequence, the intensity of X-ray diffraction peaks is not strong enough to be identified. The existence of HEA in the compound will be further confirmed hereinafter. It is worth pointing out that there is no peak belongs to the brittle secondary borides (M2 B, M23 B6 ) in the XRD patterns, which indicates that the products possess promising mechanical properties. Fig. 4 exhibits the SEM images of TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics prepared at different sintering temperatures. The secondary electron SEM images of fracture surface are shown in Fig. 4(a)–(e). Pores can be observed in Fig. 4(a), which is due to the low density of TiB2 under low temperature. With the increase of sintering temperature, the densification of TiB2 was remarkably enhanced, as shown in Fig. 4(b)–(e). However, coarser grains are also observed at the higher temperature. A widely accepted view is that ceramics materials with higher density and smaller grain size have better mechanical properties. So it is expected to achieve an optimized sintering process in this sintering system. From the fracture surface of the samples, it can be seen that fracture occurred mainly by the intergranular mode at 1300 ◦ C as a result of the high porosity. With the increase of temperature, the fracture mode evolves to the combination of transgranular and intergranular. The transgranular fracture mode may result from the strong bonding between TiB2 and HEA. A high magnification backscattered electron SEM image of 1600 ◦ C sintered TiB2 –HEA ceramics was shown in Fig. 4(f). Compared with the secondary electron SEM images, a light phase in liquid structure can be observed in the grain boundaries, which can be presumed as HEA preliminarily. It can be seen that the shape of the liquid phase is dependent on the profile of adjacent TiB2 grains. Owing to the excellent fluidity at high temperature, the liquid phase could play an important role in joining the TiB2 grains and accelerating the densification of ceramics. Moreover, the morphology also indicates a good wettability between TiB2 and HEA. For further investigating the microstructure and the mechanisms of the densification, TEM analysis was performed on the TiB2 –HEA ceramic sintered at 1600 ◦ C. Fig. 5 shows a bright field TEM image with SAED patterns of the grain and EDS patterns of liquid phase. Analysis of the SAED patterns shown in Fig. 5(b) reveals that the incident beam is paralleled to the [1¯ ¯1 2 2] zone axis of the TiB2 . The EDS patterns shown in Fig. 5(c) reveal that the liquid phase only contains the elements of Co, Cr, Fe, Ni, Ti and Al, and no B element can be detected, which evidences that the liquid phase is HEA. In addition, only a small fraction of Al can be observed, which may be related to the reactions with the oxygen-rich-layer of raw TiB2 . The HRTEM images of the boundary between ceramic grains and liquid phase are shown in Fig. 6. The selected region is magnified in Fig. 6(b) with distinct observation of a narrow grain boundary between a TiB2 ceramic grain (d = 0.204 nm, [1 0 1¯ 1] zone axis) and HEA liquid phase. This kind of tight and narrow grain boundary structure with low strain energy indicates excellent physical compatibility between TiB2 and HEA. It is well known that the grain boundary between matrix and second phases has a significant effect on the properties of the
composites, especially the mechanical properties. Therefore it is predictable that excellent mechanical properties can be obtained for the cermets. Moreover, some dispersed nanoparticles are observed to be embedded in the amorphous matrix, which is similar with the typical morphology of HEAs.21 The existence of amorphous phase can be explained by two points. First, amorphous phase would be produced during mechanical alloying. As we known, the MA processing is in favor of amorphization. Second, according to Ref. [14], there are four core effects for HEAs, which are: (1) high entropy effect; (2) sluggish diffusion effect; (3) severe lattice distortion effect and (4) cocktail effects. The effects of (1), (2), and (3) can slow down the atom diffusion rate and inhibit the crystallization. So for a HEA system with high entropy of 1.79R and severe lattice distortion, some of the amorphous phase could exist in the HEA region of the cermets. The microstructures of certain HEAs are usually very complicated, which can include nano-precipitates, ordered or disordered solid-solution phases and amorphous phases. Precipitation of nanosized particles is a novel feature of HEAs because of the effect of high mixing entropy. From the viewpoint of kinetics, long-range diffusion for phase separation is sluggish in the HEAs with multi-principal elements. Difficulties in substitutional diffusion of elements in these alloys and interactions among interdiffusing species during partitioning lower the rates of nucleation and growth, and induce the formation of nanosized crystallites. It is of great meaning and interest to investigate the coexistent of amorphous phase and nanoparticles in the liquid phase. To confirm the phase composition, the sample was further ion milled and a more identifiable HRTEM image with corresponding selected area fast Fourier transform (FFT) image of the liquid phase are shown in Fig. 7. It clearly shows that nanoparticles with size of 10 nm are embedded in the amorphous matrix and these nanoparticles are identified as a simple FCC structure by the FFT image.12 The particle size is in good agreement with that of the mechanical alloyed CoCrFeNiTiAl HEA. However, the crystal structure is different from the mechanical alloyed CoCrFeNiTiAl HEA. According to available literatures,16,22 the BCC structured HEA synthesized by mechanical alloying is a metastable state because of the non-equilibrium processing. This metastable structure will convert to a more stable phase (FCC structure here) at high temperature of annealing. That is why the HEA in the cermet is in FCC structure. In addition, FCC structure has higher atomic packing efficiency than BCC structure (FCC 74% and BCC 68%). According to Refs. [23,24], the dense packing with a large difference in the atomic size would make the alloy melt with high viscosity and low atomic mobility, which greatly reduce the nucleation rate and growth velocity, and facilitate the amorphization, even at low cooling rate. Some defects as twining crystal and lattice distortion can be observed in the nanoparticles, as shown in Fig. 7 (b)–(c), which are caused by the high energy of mechanical alloying and the large atomic radius difference between so many elements.25 This kind of structure is in favor of promoting the strength of the grain boundary phase in the cermets. The effect of the sintering temperature on mechanical properties of the TiB2 –HEA ceramics prepared by SPS is shown in
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
883
Fig. 4. The SEM images of TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics: ((a)–(e)) secondary electron images at different temperatures: (a)1300 ◦ C; (b)1400 ◦ C; (c) 1500 ◦ C; (d) 1600 ◦ C; (e) 1700 ◦ C; (f) back-scattered electron image at 1600 ◦ C.
Fig. 5. TEM images of the TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics sintered at 1600 ◦ C: (a) bright field TEM image with (b) corresponding SAED and (c) EDS patterns.
884
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
Fig. 6. The HRTEM images of the interface between ceramic grains and liquid phase for the TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics sintered at 1600 ◦ C.
Fig. 8. It reveals that the sintering temperature has a remarkable influence on the relative density, Vickers hardness and flexural strength of TiB2 –HEA ceramics. The relative density increases with increasing temperature from 1400 to 1600 ◦ C. However, further increase of sintering temperature has no positive influence on the densification of the samples. The similar trends were observed in the Vickers hardness and flexural strength, as seen in Fig. 8. The increases in hardness and strength should be mainly attributed to the improvement in bulk densities and the enhancement in bond strength between TiB2 and HEA. Grain growth restricts the further improvement of the mechanical properties at the higher temperature. The TiB2 –HEA ceramics sintered at 1600 ◦ C exhibits the attractive combination of mechanical properties, including a relative density of 99.12%,
a Vickers hardness of 2356 HV5 and a flexural strength of 800.2 MPa. A Comparison of the properties of the TiB2 -5 wt.% HEA cermets with other TiB2 -based ceramics is shown in Table 1. For monolithic TiB2 ceramics, a SPS temperature of higher than 1800 ◦ C is required for attaining high relative density (>97%).10 However, it is demonstrated by our work that the highly-dense ceramics can also be fabricated at a decreased temperature (1600 ◦ C). Besides, the TiB2 -based ceramics with a sintering aid of CoCrFeNiTiAl HEA processes a better combination of dense microstructure and excellent mechanical properties. The improvement in microstructures and mechanical properties of the ceramics in the present work due to three points: First, for most of metallic sintering aid, the formation of secondary
Fig. 7. The HRTEM images of the liquid phase for the TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics sintered at 1600 ◦ C: (a) bright field TEM image with corresponding selected area fast Fourier transform image; ((b)–(d)) magnified nano-particles in the liquid phase.
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
885
Table 1 Comparison of the properties of the TiB2 -5 wt.% HEA cermets with other TiB2 -based ceramics. Material (wt.%) Monolithic TiB2 TiB2 –5 B4 C–0.5 Fe TiB2 –2.5 Ti TiB2 –10 MoSi2 TiB2 –5 Si3 N4 TiB2 –5 AlN TiB2 –5 HEA
Processing 1800 ◦ C,
HP, 1 h, 30 MPa PS, 2000 ◦ C, 30 min SPS, 1650 ◦ C, 5 min, 50 MPa HP, 1700 ◦ C, 1 h, 30 MPa HP, 1800 ◦ C, 1 h, 30 MPa HP, 1800 ◦ C, 1 h, 30 MPa SPS, 1600 ◦ C, 5 min, 30 MPa
Relatively density (%)
Flexural strength (MPa)
Vickers hardness (GPa)
Reference
96.1 95 99.1 96.7 97.5 98 99.12
367.1 (4-P) 400 (3-P) 558 (3-P) 267.8 (3-P) 500 (3-P) 650 (4-P) 800.2 (3-P)
24 (98 N) – 26.8 (50 N) 22.97 (49 N) 22 (49 N) 21.5 (49 N) 23.56 (49 N)
[6] [11] [10] [6] [5] [4] Present work
microstructure and excellent mechanical properties, including a relative density of 99.12%, a Vickers hardness of 2356 HV5 and a flexural strength of 800.02 MPa. Acknowledgments This work was financially supported by the National Natural Science Foundation of China and the international cooperation project from Ministry of Science and Technology of China. References
Fig. 8. Mechanical properties of TiB2 -5 wt.% CoCrFeNiTiAl HEA ceramics at different sintering temperature.
borides (M2 B or M23 B6 ) embrittles the products. But in our work, no brittle phases were detected in the cermets. Second, for the crystal structure of HEA, the FCC-structured HEAs have high plasticity and relatively low yield strength, while the BCC-structured HEAs usually have very high yield strengths and limited plasticity.26 These characteristics reveal that one kind of the toughening mechanisms for this ceramic is ductile phase toughening. Third, SEM images (shown in Fig. 4) illustrate that the fracture mode changes from intergranular fracture to a combination of transgranular and intergranular fracture with the increase of relatively density. The morphologies reveal good wetting ability and bonding between the phases and also indicates crack deflection of the dense ceramic, which also promote the mechanical properties. 4. Conclusion Based on the above results and discussion, we can draw a conclusion as followed: (a) Excellent wettability (a contact angle of 16.73◦ ) between TiB2 ceramic and CoCrFeNiTiAl high-entropy alloy was detected by a high-temperature microscopy. (b) TEM observations reveal tight and narrow grain boundary between TiB2 ceramic and HEA liquid phase, which indicates excellent physical compatibility between the two components. (c) The TiB2 -based ceramic sintered at a low temperature of 1600 ◦ C provides the attractive combination of dense
1. Wen G, Li S, Zhang B, Guo Z. Reaction synthesis of TiB2 –TiC composites with enhanced toughness. Acta Mater 2001;49:1463–70. 2. Raju G, Mukhopadhyay A, Biswas K, Basu B. Densification and hightemperature mechanical properties of hot pressed TiB2 -(0–10 wt.%) MoSi2 composites. Scr Mater 2009;61:674–7. 3. Wang W, Fu Z, Wang H, Yuan R. Influence of hot pressing sintering temperature and time on microstructure and mechanical properties of TiB2 ceramics. J Eur Ceram Soc 2002;22:1045–9. 4. Li L, Kim H, Kang E. Sintering and mechanical properties of titanium diboride with aluminum nitride as a sintering aid. J Eur Ceram Soc 2002;22:973–7. 5. Park J, Koh Y, Kim H, Hwang C, Kang E. Densification and mechanical properties of titanium diboride with silicon nitride as a sintering aid. J Am Ceram Soc 1999;82:3037–42. 6. Mukhopadhyay A, Raju G, Basu B, Suri A. Correlation between phase evolution, mechanical properties and instrumented indentation response of TiB2 -based ceramics. J Eur Ceram Soc 2009;29:505–16. 7. Zheng L, Li F, Zhou Y. Preparation, microstructure, and mechanical properties of TiB2 using Ti3 AlC2 as a sintering aid. J Am Ceram Soc 2012;95:2028–34. 8. Barandika M, Sanchez J, Rojo T, Cortes R, Castro F. Fe-Ni-Ti binder phases for TiB2 -based cermets: a thermodynamic approach. Scr Mater 1998;39:1395–400. 9. Einarsrud M, Hagen E, Pettersen G, Grande T. Pressureless sintering of titanium diboride with nickel, nickel boride and iron additives. J Am Ceram Soc 1997;80:3013–20. 10. Zhang Z, Shen X, Wang F, Lee S, Fan Q, Cao M. Low-temperature densification of TiB2 ceramic by the spark plasma sintering process with Ti as a sintering aid. Scr Mater 2012;66:167–70. 11. Kang E, Jang C, Lee C, Kim C, Kim D. Effect of iron and boron carbide on the densification and mechanical properties of titanium diboride ceramics. J Am Ceram Soc 1989;72:1868–72. 12. Zhang J, Fu Z, Wang W, Wang H, Min X. Wettability between TiB2 ceramic and metals. Acta Metall Sin 1999;12:395–400. 13. Zhu G, Liu Y, Ye J. Fabrication and properties of Ti(C, N)-based cermets with multi-component AlCoCrFeNi high-entropy alloys binder. Mater Lett 2013;113:80–2. 14. Yeh J, Chen S, Lin S, Gan J, Chin T, Shun T, et al. Nanostructured highentropy alloys with multiple principal elements: novel alloy design concepts and outcomes. Adv Eng Mater 2004;6:299–303.
886
W. Ji et al. / Journal of the European Ceramic Society 35 (2015) 879–886
15. Otto F, Yang Y, Bei H, George E. Relative effects of enthalpy and entropy on the phase stability of equiatomic high-entropy alloys. Acta Mater 2013;61:2628–38. 16. Ji W, Fu Z, Wang W, Wang H, Zhang J, Wang Y, et al. Mechanical alloying synthesis and spark plasma sintering consolidation of CoCrFeNiAl highentropy alloy. J Alloy Compd 2014;589:61–6. 17. Zheng X, Shen P, Han X, Lin Q, Qiu F, Zhang Y, et al. Wettability and reactivity between B4 C and Molten Zr55 Cu30 Al10 Ni5 metallic glass alloy. Mater Chem Phys 2009;117:377–83. 18. Yeh J, Chang S, Hong Y, Chen S, Lin S. Anomalous decrease in X-ray diffraction intensities of Cu–Ni–Al–Co–Cr–Fe–Si alloy systems with multiprincipal elements. Mater Chem Phys 2007;103:41–6. 19. Aizenshtein M, Froumin N, Frage N. The nature of TiB2 wetting by Cu and Au. J Mater Eng Perform 2012;21:655–9. 20. Varalakshmi S, Kamaraj M, Murty B. Synthesis and characterization of nanocrystalline AlFeTiCrZnCu high entropy solid solution by mechanical alloying. J Alloy Compd 2008;460:253–7.
21. Zhang X, Zhang Y, Qiao Y, Chen G. Novel microstructure and properties of multicomponent CoCrCuFeNiTix alloys. Intermetallics 2007;15: 357–62. 22. Ji W, Wang W, Wang H, Zhang J, Wang Y, Zhang F, et al. Alloying behavior and novel properties of CoCrFeNiMn high-entropy alloy fabricated by mechanical alloying and spark plasma sintering. Intermetallics 2015;56:24–7. 23. Li Y, Guo Q, Kalb, Thompson C. Matching glass-forming ability with the density of the amorphous phase. Science 2008;322:1816–9. 24. Li R, Pang S, Men H, Ma C, Zhang T. Formation and mechanical properties of (Ce–La–Pr–Nd)–Co–Al bulk glassy alloys with superior glass-forming ability. Scr Mater 2006;54:1123–6. 25. Fu Z, Chen W, Fang S, Zhang D, Xiao H, Zhu D. Alloying behavior and deformation twinning in a CoNiFeCrAl0.6 Ti0.4 high entropy alloy processed by spark plasma sintering. J Alloy Compd 2013;553:316–23. 26. Zhang Y, Zuo T, Tang Z, Gao M, Dahmen K, Liaw P, et al. Microstructures and properties of high-entropy alloys. Prog Mater Sci 2014;61:1–93.