Microstructural evolution and mechanical properties of TiB2-TiC-SiC ceramics joint brazed using Ti-Ni composite foils

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

Original Article

Microstructural evolution and mechanical properties of TiB2-TiC-SiC ceramics joint brazed using Ti-Ni composite foils X.Q. Caia,b, D.P. Wanga,b, Y. Wanga,b,*, Z.W. Yanga,b a

Tianjin Key Lab of Advanced Joining Technology, School of Materials Science and Engineering, Tianjin University, Tianjin, 300072, China Key Laboratory of Advanced Ceramics and Machining Technology of Ministry of Education, School of Materials Science and Engineering, Tianjin University, Tianjin, 300072, China

b

ARTICLE INFO

ABSTRACT

Keywords: TiB2-TiC-SiC ceramics Brazing TiB whisker Microstructural evolution Mechanical properties

A novel TiB2-based ultra-high-temperature ceramic containing 60 vol.% TiB2, 20 vol.% TiC, and 20 vol.% SiC was fabricated by hot pressing and subsequently joined using the brazing technique. Ti-based filler was used as the brazing alloy by taking advantage of the reaction between Ti and TiB2-TiC-SiC. The effects of the brazing temperature on the microstructure and mechanical properties of the brazed joint were investigated. The results showed that Ti in the filler reacted with the TiB2-TiC-SiC ceramics and formed a reaction layer I that comprised TiB and TiC. The brazing seam was composed of TiB, TiC, Ti5Si3, Ti2Ni, and TiNi. When the brazing temperature was increased, the reaction between TiB2-TiC-SiC ceramics and the filler was observed to become vigorous; this led to an increase in the growth of the reaction layer I. Meanwhile, the continuous Ti2Ni layer in the brazing seam gradually disappeared; it was replaced by TiB and Ti5Si3. The room temperature shear strength reached a maximum value of 168 MPa when the joint was brazed at 1040 °C for 30 min; while it was 104 and 81 MPa at test temperature of 600 °C and 800 °C, respectively. In addition, the effects of TiB whiskers on the coefficient of thermal expansion of the brazing seam and fracture of the brazed joint were discussed.

1. Introduction Ultra-high-temperature ceramics (UHTCs), especially transition metal diborides, have recently received increased attention owing to their broad range of technological applications [1–4]. Further, titanium diboride (TiB2)-based UHTCs, in particular, have received even more attention due to the attractive combinations of their excellent thermophysical, chemical and mechanical properties [5–7]. Compared with other transition metal diborides ceramics, TiB2 ceramics have the lowest density and coefficient of thermal expansion (CTE) [1]; this is advantageous for applications that require a higher thrust-to-weight ratio. Although TiB2 ceramics possess excellent properties, monolithic TiB2 ceramics are difficult to densify and they have very low fracture toughness [8,9]. Moreover, Tampieri et al. revealed that monolithic TiB2 ceramics underwent extensive oxidation at temperatures above 1000 °C [10]. To solve these problems, various non-metallic additives have been utilized to enhance the densification and mechanical properties of TiB2-based UHTCs [11–14]. Among these additives, titanium carbide (TiC) shows appropriate thermodynamic compatibility and coherency with TiB2 [15]; it is commonly used to improve the

densification and fracture toughness of TiB2-based UHTCs [11,15]. Silicon carbide (SiC), on the other hand, is considered as an efficient additive to improve the oxidation resistance of these ceramics [13,16]. Thus, TiB2-TiC-SiC ceramics possess a combination of excellent properties that are shared by each component [17,18]. That will broaden the scope of use for these ceramics. However, engineering applications require an assembly of components for the fabrication of large and complex parts. Therefore, joining the TiB2-based UHTCs is indispensable to the industrial applications. To date, fusion welding [19] techniques have been employed to join TiB2-based UHTC. Derek et al. investigated plasma arc welding of TiB220TiC (vol.%) ceramics [19]. Although the results were satisfactory, the formation of large crystals and pores in the fusion zone were unavoidable during the cooling cycle. The brazing technique is an economical and feasible joining method, and has been gaining attention as the best alternative to join UHTCs [20–26]. Ag-based [20,21], Ni-based [22,23], Ti-based [24,25] and Pa-based [26] interlayers have been applied to join UHTCs. Even though very few reports have actually addressed its use in literature about the joining of TiB2-based UHTCs. A key problem with the UHTCs brazing method, however, has been the

⁎ Corresponding author at: Tianjin Key Lab of Advanced Joining Technology, School of Materials Science and Engineering, Tianjin University, Tianjin, 300072, China. E-mail addresses: [email protected], [email protected] (Y. Wang).

https://doi.org/10.1016/j.jeurceramsoc.2020.03.053 Received 12 November 2019; Received in revised form 27 February 2020; Accepted 24 March 2020 0955-2219/ © 2020 Published by Elsevier Ltd.

Please cite this article as: X.Q. Cai, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2020.03.053

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

residual stress that are attributable to a mismatch in the CTE between ceramic and brazing alloy. Because of good compatibility between TiB2 and TiB [27,28], the formation of TiB in TiB2-based UHTCs joint is an excellent choice; further, TiB can be obtained when Ti and TiB2 react. Designing the brazing alloy with the active element of Ti is crucial as it renders the TiB2-based UHTCs joint more reliable. AgCu + Ti [20,21,29], Ti-Cu-Ni [30], Ti-Zr-Ni-Cu [31,32], and Ti-Ni [24,25] brazing alloys with the active element of Ti are potential fillers to join UHTCs, and appropriating joint performance could be obtained [24,25]. For the practical application of UHTCs, high-temperature mechanical performance of the joint also emerges as a problem. The mechanical stability of the joint brazed with the Ag-based brazing alloy was found to dramatically degrade above 450 °C [33]. The addition of Cu was beneficial to reduce the brazing temperature of the Ti-basedbrazing alloy. However, the formation of Ti-Cu intermetallics in the joint restricted the high-temperature performance of the joint. Among these brazing alloys that comprise an active element of Ti, Ti-Ni alloy was the best fit to be utilized in high-temperature environments. The TiNi brazing alloy was used to braze ZrB2-SiC ceramics, and a sound joint with a maximum shear strength of 134 MPa at room temperature was obtained [25]. Based on the above analysis, Ti-Ni brazing alloy may be a good choice for the brazing of TiB2-based UHTCs. In this study, a novel TiB2-based UHTC containing 60 vol.% TiB2, 20 vol.% TiC, and 20 vol.% SiC was fabricated by hot pressing; further, Ti-Ni filler was used as the brazing for TiB2-TiC-SiC ceramics. The interfacial microstructure and evolution mechanism of the brazed joints were analysed in detail. The effects of the brazing temperature on the microstructure and mechanical properties of the brazed joint were also investigated. In addition, the effects of the formation of TiB whiskers on the coefficient of thermal expansion of the brazing seam, and the fracture of the brazed joint were discussed.

speed = 0.05 mm/min). These values conformed to the standard (ISO 14704: 2000). The flexural strength of TTS ceramic was conducted using a universal testing machine (MTS Model E45.106). The flexural strength of ceramic was averaged from five specimens. 2.3. Brazing process The sizes of TTS ceramic brazing specimens were 5 mm × 5 mm × 5 mm and 15 mm × 10 mm × 3 mm, respectively. The joining surfaces of the TTS ceramic were polished using diamond millstones up to 3000 grit, and then, the polished TTS ceramic was ultrasonically cleaned in an acetone solution about 10 min. The commercially available pure Ti and Ni foils were used in this experiment. Based on the lowest eutectic point (942 °C) in Ti-Ni binary phase diagram [34], the assembly structure of TTS / Ti (60 μm) / Ni (25 μm) / Ti (60 μm) / TTS was designed in this study. To keep a proper contact of all the parts, a pressure of 2 kPa was applied to the assemblies. Brazing experiments were performed in a vacuum furnace (ZT-40-21Y, Shanghai Chen Hua Electric Furnace Co., Ltd., China). Prior to heating, the vacuum furnace was evacuated to a pressure of 1.0 × 10−3 Pa. Chen et al. [35] pointed out that the brazing temperature was generally at least 60 °C above the Ti-Ni liquidus temperature to ensure that the filler alloy was entirely melted during brazing. Therefore, the brazing temperatures ranging from 1000 °C to 1060 °C are selected in this study. The heating rate was 10 °C/min. The holding time was 30 min. The cooling rate was 10 °C/min between peak brazing temperature and 400 °C, and then followed cool to room temperature in the furnace. 2.4. Characterisation of brazed joint After brazing, cross-sections of the TTS ceramic brazed joints were wet-ground and polished. The microstructures of the joints were characterized by scanning electron microscopy (SEM). Chemical composition analysis of the reaction products in the joints was performed using an energy-dispersive spectroscopy (EDS) apparatus with different acceleration voltage ranging from 15 to 20 kV. A standardless quantitative analysis was performed employing the ZAF correction approach. The reaction phases were further observed by transmission electron microscopy (TEM, JEM-ARM200F). The lattice structure of the reaction phases was determined by selected area electron diffraction (SAED). Specific samples for TEM observation were prepared using a focused ion beam system (FEI Helios NanoLab 600i). Shear tests were conducted using a universal testing machine (MTS Model E45.106) to evaluate the joint strength. The brazed specimen was compressed with a constant speed of 0.1 mm/min. The shear strength of brazed joint was measured at room temperature (RT) and at elevated temperatures of 600 °C and 800 °C, respectively. At least four specimens were used for each experimental condition to average the joint strength. After the shear test, fracture paths and fracture morphologies of the brazed joints were characterized by SEM. The reaction phase of the fracture surface was identified using X-ray diffractometer (XRD). Nano-indentation measurements were conducted using a Nano Indenter XP equipped with a diamond Berkovich tip. Indentations were conducted on the polished cross-section of TTS ceramics brazed joint. The load was applied to a maximum depth of 500 nm. Values of hardness (H) and Young's modulus (E) of reaction phases were obtained from the load-displacement data for the indentations using the Oliver-Pharr method [36]. At least three positions were used for each phase to ensure reliability of the data.

2. Materials and experimental procedures 2.1. Fabrication of TiB2-TiC-SiC ceramics Commercial powders were used to prepare the TiB2-TiC-SiC ceramics. The TiB2 powder (purity: 99.5%, Beijing Huawei Ruike Chemical Co., Ltd, Beijing, China) and TiC powder (purity: 99.5%, Beijing Huawei Ruike Chemical Co., Ltd, Beijing, China) had particle sizes of 1–2 μm. The SiC powder (purity: 99.0%, Shanghai Macklin Biochemical Co., Ltd, Shanghai, China) was predominantly β-SiC, and it had particle sizes that ranged from 0.5 to 0.7 μm. The ceramic with a nominal composition of 60 vol.% TiB2, 20 vol.% TiC, and 20 vol.% SiC was fabricated by a hot pressing technique. The material is referred to as TTS ceramic hereafter. The raw powders were mixed in alcohol by the wet-ball milling for 24 h in a polyethylene bottle using tungsten carbide (WC) balls. The powder mixtures were then dried in a dry vacuum evaporator and thereafter sieved through a 200-mesh sieve for further use. The dried powder was sintered via hot pressing type at 2000 °C in a graphite die under a vacuum atmosphere with a holding time of 60 min at 30 MPa of pressure. The sintered specimens were in the form of pellets with 50 mm diameter and approximately 6 mm thickness. 2.2. Characterisation of TiB2-TiC-SiC ceramics The density of the TTS ceramic was measured by Archimedes’ method in distilled water. Relative density was calculated using a volumetric rule of mixtures, assuming the true densities: 4.52 g/cm3 for TiB2, 4.93 g/cm3 for TiC, and 3.2 g/cm3 for SiC. The microstructure of the TTS ceramic was characterized by scanning electron microscopy (SEM, JSM-7800 F). The phase composition of the TTS ceramic was observed using an X-ray diffractometer (XRD, Bruker D8 Advanced) using Cu Kα radiation over ranging from 20 to 90°. The flexural strength of ceramic was detected by three-point bending tests (specimen size = 30 mm × 5 mm × 5 mm, span = 20 mm, loading

3. Results and discussion 3.1. Characterisation of TiB2-TiC-SiC ceramics The density of the hot-pressed TTS ceramic was 4.26 g/cm3, corresponding relative density of 98%. The flexural strength of the TTS 2

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 1. Microstructure (a) and the XRD pattern (b) of the TiB2-TiC-SiC ceramic.

ceramic was 664 ± 32 MPa. Fig. 1 shows the backscattered scanning electron (BSE) micrograph and X-ray diffraction (XRD) pattern of the hot-pressed TTS ceramic. The results show that TTS ceramic primarily consisted of greyish TiB2 matrix, off-white TiC, and dark grey SiC; the TiB2, TiC, and SiC were uniformly distributed.

Table 1 Average chemical composition analysis of each phase in Fig. 2 (at.%).

3.2. The typical interfacial microstructure of TTS/Ti-Ni/TTS brazed joint Fig. 2 shows the typical backscattered electron (BSE) images of the TTS ceramic joint brazed using Ti-Ni filler at 1040 °C for 30 min. A sound joint was obtained by using Ti-Ni filler; cracks or pores were not detected in the brazed joint. Moreover, a large number of new phases were formed in the brazed joint, indicating that chemical reactions occurred between the liquid Ti-Ni filler and TTS ceramic during brazing. The joint mainly consisted of two reaction zones: Zone I (reaction layer adjacent to TTS ceramic) and Zone II (central brazing seam). The average chemical composition of reaction phases was measured by EDS, and the results are listed in Table 1. Fig. 2 (b) shows the magnified microstructure of the reaction layer I. This mainly consist of the A phase. According to the composition analysis in Table 1, the reaction layer I consisted of 46.24 at.% boron and 50.57 at.% titanium, and the stoichiometric ratio of titanium and boron in this phase indicated a TiB phase. In addition, the B phase bordering around the SiC was also detected in the reaction layer I. This reaction phase primarily consisted of 48.84 at.% titanium, 36.34 at.% carbon. According to the TiC phase diagram [37], phase B bordering around SiC was TiC. Fig. 2(c) shows the magnified microstructure of the brazing seam according to which the brazing seam consisting of five different phases: C, D, E, F and G. Needle-shaped phase C primarily consisted of titanium and boron with stoichiometric ratio of 1:1. This phase was considered as TiB. The crystal structure and formation mechanism of TiB whiskers will be discussed later. Fig. 3 shows the bright-field TEM image of the reaction phases formed in the brazing seam and the corresponding selected area

Position

Ti

Ni

B

C

Si

Possible phase

A B C D E F G

50.57 48.84 49.29 54.16 50.16 53.84 44.12

00.54 00.43 00.55 04.16 08.62 24.34 41.35

46.24 10.63 47.50 13.82 09.83 08.87 07.66

02.26 36.34 02.28 05.63 30.35 11.06 06.06

00.39 03.76 00.38 22.23 01.04 01.89 00.81

TiB TiC TiB Ti5Si3 TiC Ti2Ni TiNi

electron diffraction (SAED) patterns. According to these patterns and the composition results in Table 1, phase D mainly consisted of 54.16 at.% titanium, and 22.23 at.% silicon; it was considered as Ti5Si3. Phase E mainly consisted of titanium and carbon, which was regarded as TiC. The titanium-nickel ratio of phase F is about 2:1, and it was considered as Ti2Ni. Conversely, phase G was not detected in the TEM results; it also consisted of titanium and nickel. The titanium-nickel ratio in phase G was about 1:1. According to the Ti-Ni phase diagram [34], phase G was TiNi. XRD analysis on the fracture surface of the brazed joint is shown in Fig. 4, containing the formation of TiB, TiC, Ti5Si3, Ti2Ni, and TiNi phases in the brazed joint. In conclusion, the interfacial structure of TTS/ Ti-Ni/TTS joint brazed at 1040 °C for 30 min is: TTS/TiB + TiC/ TiB + TiC + Ti5Si3+Ti2Ni + TiNi. Based on the above interfacial microstructure analysis, the reaction layer I adjacent to TTS ceramic was mainly occupied by TiB, and a large number of TiB whiskers were formed in the brazing seam. Thus, the formation of TiB was the key factor to realise the joining of TTS ceramic. TEM analyses were performed to observe morphologies of the reaction layer as well as the TiB structure. Fig. 5 shows the bright-field TEM image of the reaction layer adjacent to TiB2 component and the corresponding SAED patterns. As observed in Fig. 5(a), the reaction layer adjacent to the TiB2 component was composed of TiB, indicating

Fig. 2. Microstructure of TTS/Ti-Ni/TTS joint brazed at 1040 °C for 30 min: (a) integral joint, (b) magnification of the reaction layer I, and (c) magnification of the brazing seam. 3

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 3. TEM analysis of the brazing seam: (a) bright-field image, and corresponding selected area electron diffraction (SAED) pattern of (b) Ti5Si3, (c) TiC, and (d) Ti2Ni.

The crystal structure of the TiB was B27 structure [38,39]; it consisted of trigonal prisms stacked in columnar arrays that shared only two of their three rectangular faces with adjacent prisms. The B atom was located at the centre of the triangular prism of the six Ti atoms, and therefore, the B atom formed a zigzag chain along the [010] direction. The crystal structure of TiB had a very high asymmetry, indicating that TiB would grow in anisotropy. In the (100), (001) and (010) planes of TiB, the arrangements of Ti and B atoms were different. In the (100)TiB plane, Ti and B atoms occupied the sites of the alternating layers [39]. In the (010)TiB and (001)TiB planes, the atom ratio between Ti and B was 1:1. The normal growth in planes containing both the Ti and B atoms in the same stoichiometry as the crystal should be faster than the growth along the directions involving alternating planes of Ti and B atoms [40]. Thus, the growth to the (100)TiB plane should be slower than that of (010)TiB and (001)TiB planes. When TiB nucleated and grew from the liquid filler, the (100)TiB plane retained in the joint because the growing velocity of the (100) plane was slower than that of the (001) plane, and the (001) plane disappeared in the transverse section of TiB. Therefore, TiB whisker had a hexagonal shape in the transverse section, and it consisted of three pairs of low-index crystal planes of (100), (101) and (101). The formation of stacking faults in TiB whiskers was dependent on the B atoms in the liquid filler. The number of Ti atoms in the liquid filler was sufficient. The number of B atoms in the liquid filler was significantly less than that of Ti atoms. When there were lack of B atoms, stacking faults would form. Stacking faults formed in the (100)TiB planes had minimum energy compared with those formed in the (001)TiB and (010)TiB planes [40]. Moreover, The stacking sequence along the [010]TiB and [001]TiB directions was more simple than that along the [100]TiB direction [39,41]. Accordingly, stacking faults in (100)TiB formed in TiB whiskers during the brazing process. Based on the above analysis, the microstructural evolution mechanism of TTS ceramic joints brazed with Ti-Ni composite foils was divided into the following four stages.

Fig. 4. XRD pattern of the fracture surface of TTS ceramic joint brazed at 1040 °C for 30 min.

the reaction between Ti in the filler and TiB2 ceramic was complete. TiB phases first nucleated and then grew along the surface of the TiB2 ceramic. TiB grains were joined to each other; as a result, they formed a continuous TiB layer. Fig. 6 shows the bright-field TEM image of the TiB whisker formed in the brazing seam along the transverse section and the corresponding SAED pattern. The SAED pattern of TiB whisker indicated that the crystallographic planes of the transverse section of TiB were always along (100), (101) and (101) planes. The TiB whisker formed in the brazing seam exhibited a growth orientation in the [010]TiB direction. Fig. 7 show the bright field TEM image of TiB whisker formed in the brazing seam along the longitudinal section and the corresponding SAED pattern, and the high-resolution TEM image of stacking faults in the TiB whisker. The results revealed that a large number of stacking faults were present in TiB whiskers. The SAED and high-resolution TEM patterns showed that the stacking faults were always parallel to the (100)TiB plane.

(1) First stage (physical contact): According to the Ti-Ni binary phase diagram [34], the α-Ti moved to β-Ti when the temperature increased to 882 °C. In addition, the Ti and Ni foils were closely 4

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 5. TEM analysis of the reaction layer adjacent to the TiB2 component: (a) bright-field image with corresponding SAED patterns of (b) TiB2, and (c) TiB.

joined, and Ti and Ni atoms migrated by solid-phase diffusion. The intermetallic compound of Ti2Ni formed between the Ti and Ni foils during heating. (2) Second stage (melting and diffusion of atoms): As the temperature increased to 942 °C, the β-Ti-Ti2Ni eutectic point was reached. The eutectic liquid L formed according to the reaction path β-Ti + Ti2Ni →L. Ti atoms in the liquid phase diffused and aggregated near the TTS ceramics surface, driven by the gradient of concentration between the Ti-Ni liquid filler and TTS ceramics. (3) Third stage (forming reaction products): In order to further analyse the reaction products of the TTS ceramic and Ti-Ni liquid filler, the details of the TiB2, SiC and TiC interface are shown in Fig. 8(a)–(c). The compositions of each phase marked in Fig. 8 were measured by EDS, and the results are listed in Table 2.

filler reacted with TiB2 ceramic to form TiB; the reaction could be proceed following Reaction (1). The formation of TiB in the brazed joint 0 was thermodynamically feasible ( G1313 K ＜0). Owing to the sufficient Ti in the Ti-Ni liquid phase, the reaction would proceed continuously until a certain thickness of TiB reaction layer. During the formation of a certain thickness of TiB reaction layer, part of TiB migrated into the central of Ti-Ni liquid phase. Therefore, TiB whiskers were formed in the brazing seam, as shown in Fig. 2 (c). Ti + TiB2 → TiB ΔG0 (kJ/mol) = −41.946 + 0.0825 T.

(1)

On the interface of the SiC component, Ti-Ni liquid phase was in contact with the SiC ceramic. The active element Ti reacted with SiC ceramic and yielded TiC and Si, according to Reaction (2). A further reaction between SiC ceramic and Ti would be hindered since the previous TiC layer impeded Ti diffusion to SiC ceramic. As a result, a

On the interface of the TiB2 component, the Ti in the molten Ti-Ni

Fig. 6. TEM analysis of the TiB whisker formed in the brazing seam along the transverse section: (a) bright-field image, and (b) corresponding SAED pattern. 5

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 7. TEM analysis of TiB whisker formed in the brazing seam along the longitudinal section: (a) bright-field image, (b) corresponding SAED pattern, and (c) TEM image of stacking faults in the TiB whisker.

thin TiC layer formed adjacent to the SiC ceramic, as shown in Fig. 8(b). During the formation of the TiC reaction layer, part of Si dissociated and diffused into the Ti-Ni liquid phase to react with Ti and form Ti5Si3 according to Reaction (3). Therefore, Ti5Si3 phases formed in the brazing seam, as shown in Fig. 2(c). The Gibbs free energy of reaction (2) and (3) was negative at 1313 K. Thermodynamic calculation from the free energy of TiC and Ti5Si3 formations was feasible. The interfacial microstructure of the SiC ceramic brazed with filler containing Ti has been extensively studied [42,43]. Furthermore, a thin TiC layer near SiC ceramic and Ti5Si3 phase in the brazing seam was formed. Ti + SiC → TiC + Si ΔG° (kJ/mol) = −136.9 + 0.0095T

(2)

5/3Ti + Si → 1/3Ti5Si3 ΔG0 (KJ/mol) = −194.14 + 0.0167T

(3)

Table 2 Average chemical composition analysis of each phase in Fig. 8 (at. %). Ti

Ni

B

C

Si

Possible phase

A B C D E

50.11 43.59 48.17 58.78 43.35

01.94 01.14 00.67 01.00 00.89

42.60 11.67 43.77 07.18 08.02

04.06 35.65 06.54 32.07 42.82

01.29 07.96 00.85 00.97 03.92

TiB TiC TiB TiC TiC

(4) Fourth stage (forming the joint): During cooling, Ti2Ni precipitated in the residual liquid phase as the temperature decreased to its melting point. In addition, TiNi also precipitated in the brazing seam due to Ti deficience in some regions.

On the interface of TiC component, the reaction between Ti-Ni liquid phase and TiC ceramic can be divided into two stages. Firstly, the interdiffusion would occur when TiC ceramic is directly in contact with solid Ti in a physical contact stage. Secondly, Ti in the liquid phase could promote the interdiffusion because of its fast diffusion rate in the Ti-Ni liquid phase during brazing. Ti from the liquid filler diffused into TiC ceramic and C atoms from TiC ceramic diffused into the liquid filler. Thus, TiC ceramic with lower C/Ti ratio was found in the TTS ceramics adjacent to Ti-Ni filler interface as shown in position D of Fig. 8(c). However, the C/Ti ratio of TiC ceramic in the TTS ceramic was close to 1 as shown in position E of Fig. 8(c). As a non-stoichiometric compound, the ratio of Ti to C was not fixed in the TiC crystal. TiC exists as a single phase for C/Ti ratio ranging from 0.48 to 1 [29,37]. In addition, C diffused into the Ti-Ni liquid phase and reacted with Ti to form TiC phase in the brazing seam according to the following reaction: C + Ti → TiC ΔG0 (KJ/mol) = −186.6 + 0.0132 T

Position

3.3. Effect of brazing temperatures on the interfacial microstructure of TTS/ Ti-Ni/TTS brazed joints Fig. 9 shows the interfacial microstructures of TTS/Ti-Ni/TTS joints brazed at different temperatures for 30 min. The interfacial structure and reaction products revealed some differences. With an increase in the brazing temperature, the thickness of the reaction layer I increased, as observable in Fig. 9(a)–(d). A continuous Ti2Ni phase area appeared in the brazing seam when the brazing temperature was low (e.g. 1000 °C and 1020 °C). Whereas was absent with an increase in the brazing temperature, substituted by a large number of Ti5Si3 and TiB particles. In addition, the size of Ti5Si3 and TiB phases in the brazing seam increased at the elevated temperature. Another interesting phenomenon was that a certain amount of off-white TiNi phases formed in the brazing seam at 1040 °C and 1060 °C, as seen in Fig. 9(c) and (d). TiNi phase could not be found when the brazing temperature was lower, as observed in Fig. 9(a) and (b). The change of the interfacial structure and reaction products depends on the reaction between the TTS ceramic and Ti-Ni filler during the brazing process. The activity of Ti was gradually improved with

(4)

During the brazing process, TiB nucleated and grew along the surface of TTS ceramic, and finally occupied the interface of the TiC ceramic. Thus, the interface of TiC ceramic was occupied by TiB, as shown in Fig. 8(c).

Fig. 8. Microstructure of TTS/Ti-Ni/TTS joint brazed at 1040 °C for 30 min: (a) detail of TiB2 interface, (b) detail of SiC interface, and (c) detail of TiC interface. 6

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 9. Microstructure of TTS/Ti-Ni/TTS joints brazed for 30 min at different temperatures: (a) 1000 °C, (b) 1020 °C, (c) 1040 °C, and (d) 1060 °C.

increasing brazing temperature accompanied by increasing amount of Ti-Ni liquid phase. With an increase in brazing temperature, the reaction between TiB2 and Ti became vigorous, and TiB increasingly formed in the reaction layer I. Thus, the thickness of the reaction layer I increased with an increase in the brazing temperature. On the other hand, an increase in the brazing temperature could contribute to TiB migrated into the central of Ti-Ni liquid phase; this promoted the formation of TiB whiskers in the whole brazing seam. However, the brazing temperatures were not the highest or the best for the formation of TiB whiskers. TiB whiskers became coarse, especially when the brazing temperature was 1060 °C. Similarly, the reaction between SiC and Ti became vigorous at a high temperature, and increasingly Si diffused into the Ti-Ni liquid phase. Ti5Si3 phase nucleated in the whole brazing seam and gradually grew. Finally, the coarse Ti5Si3 phase formed in the brazing seam at elevated brazing temperatures. The formation of TiB and Ti5Si3 phases in the brazing seam consumed a mass of Ti atoms in the liquid phase at high temperature. Lack of Ti atoms was likely to happen in some areas. Thus, there were some TiNi phases formed in residual liquid phase at higher brazing temperatures.

Table 3 Hardness and Young’ modulus of the reaction phases in the TTS/Ti-Ni/TTS joint brazed at 1040 °C for 30 min. Phase

TiB2 TiBT TiBL Ti2Ni Ti5Si3 TiC

Average

Typical

Hardness/GPa

Modulus/GPa

Hardness/GPa

Modulus/GPa

31.2 20.7 20.5 12.2 17.4 22.9

510 385 351 198 280 355

31.3 20.9 20.6 12.6 17.2 22.6

520 390 341 194 285 351

± ± ± ± ± ±

0.5 0.6 1.0 1.5 0.8 0.6

± ± ± ± ± ±

16 40 31 26 20 18

the brazing seam at transverse and longitudinal sections were also indented. For convenience, TiB whiskers at transverse and longitudinal sections were marked as TiBT and TiBL, respectively. The results showed that the TiB2 in TTS ceramic had the highest E and H of ∼520 GPa and ∼31 GPa, respectively, which was significantly higher than the reaction phases formed in the brazing seam. For the same penetration depth, the maximum load of the indenter for TiB2 was approximately 2.3 times more than that measured for the Ti2Ni phase. The Ti2Ni phase had the lowest E and H of about 194 GPa and 13 GPa, respectively. In contrast, Ti5Si3 and TiC were harder than Ti2Ni. Their parameters were as follows: the E and H of about 285 GPa and 17 GPa for Ti5Si3, respectively and about 351 GPa and 23 GPa for TiC, respectively. Additionally, the hardness of TiB at the transverse section was similar to that of TiB at the longitudinal section. The hardness of TiB at transverse and longitudinal section was about 21 GPa. To evaluate the deformation behaviours of the reaction phases, the typical load-displacement curves of these phases are illustrated in Fig. 10(b). According to the report [44], plastic factor η of TiB2 can be obtained as:

3.4. Mechanical properties and fracture analysis of TTS ceramics brazed joints The nature of the reaction phases is of paramount importance for the mechanical properties of the joint. To better focus on the nature of the reaction phases, nano-indentation was performed on the cross-sections of the TTS ceramic brazed joint. The hardness (H) and Young's modulus (E) of the reaction phases are listed in Table 3. Fig. 10 shows the H and E of the reaction phases and the typical load-displacement curves of these phases. SEM in BSE mode was adopted for assurance of the indentation positions, as shown in Fig. 10(c)–(h). All reaction phases except TiNi were indented. In addition, TiB whiskers formed in 7

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 10. (a) Hardness and Young’ modulus of the reaction phases in the TTS/Ti-Ni/TTS joint brazed at 1040 °C for 30 min, (b) load-displacement curves for the reaction phases, and the corresponding indentation positions of reaction phases: (c) TiB2, (d) TiBT (TiB along the transverse section), (e) TiBL (TiB along the longitudinal section), (f) Ti2Ni, (g) Ti5Si3, and (h) TiC.

η(TiB2)= SABC / SABD

was 1000 °C. According to Dragan Toprek research, CTE of Ti2Ni reached 40.48 × 10−6 K−1 at 300 K [45]. The CTE of TTS ceramic was conducted using a thermal diatomter (NETZSCH DIL 402CL) according to the standard (GB/T 4339). The CTE of TTS ceramic was 5.58 × 10−6 K-1 at 300 K. Compared with TTS ceramic, the CTE of Ti2Ni phase was far more than that of TTS ceramic. Large CTE mismatch would lead to large residual stress in the brazed joint, which harm the joint strength. Higher brazing temperatures promoted the formation of TiB whiskers in the brazing seam and the TiB whiskers in the brazing seam could accommodate the residual stress. The brazing seam mainly consisted of Ti2Ni phase and TiB whiskers, which could be regarded as a Ti2Ni composite reinforced by TiB whiskers. CTE of the composite was relatively difficult to predict because it is dependent on several factors such as the internal structure of the composite and the plasticity of the matrix. Existing literature presents the CTE of composites predicated by different classical models [46–48]. The classical Rom model [46], Turner model [47] and Kerner’s model [48] were the simplest and most often used, and corresponding the mathematical presentations of these three models are listed in Eqs. (5)–(7) as shown below:

where, SABC and SABD are the areas of ABC and ABD in Fig. 10(b), respectively. From the calculation, the plastic factor η of TiB2 was 51.9%. The plastic factors η of TiBT, TiBL, Ti2Ni, Ti5Si3, and TiC were 61.6, 62.0, 66.2, 63.6 and 59.8, respectively. Compared with the TiB2 in TTS ceramic, the reaction phases formed in the joint had a better plastic deformation behaviour. According to the plastic factor η, Ti2Ni had good plasticity. We suppose that the plastic deformation could absorb the part of the detrimental residual stress generated in the brazed joint by the mismatch of the coefficient of thermal expansion. Fig. 11(a) shows the effect of brazing temperature on the average room temperature (RT) shear strength of TTS ceramic brazed joints. The shear strengths of the joints increased and then decreased with an increase in the brazing temperature. The maximum RT shear strength of 168 MPa was achieved when the joint was brazed at 1040 °C for 30 min. TTS ceramic is usually applied to high-temperature environments. Results of, it is necessary to investigate the high-temperature behaviour of the brazed joint; the high temperature (600 °C and 800 °C) shear tests are shown Fig. 11(b). As observed, the shear strength of the joints decreased dramatically with an increase in the test temperature. The shear strengths of the joint were 104 and 81 MPa at 600 °C and 800 °C, respectively. The shear strength depended on the interfacial microstructure of TTS ceramic-brazed joints. As discussed above, a continuous Ti2Ni phase area appeared in the brazing seam when the brazing temperature 8

1

=

2

=

3

=

w Vw

(5)

+ w Kw Vw KmVm + Kw Vw

(6)

m Vm

+

m Km Vm

m Vm

+

w Vw

+

Vm Vw (

m

w )(Km

Vm Km + Vw Kw +

(

Kw ) 3KmKw 4G w

)

(7)

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 11. Effect of (a) the brazing temperature and (b) the test temperature on the shear strengths of TTS ceramics brazed joints.

where the subscripts m and w indicate the Ti2Ni matrix and TiB whiskers, respectively; α is the CTE value; V is the volume fraction; K is the bulk modulus, and G is the shear modulus. In Rom’s model, the CTE of the composite is considered as the weighted sum of the CTE of each component. Meanwhile, the interphase interaction between the Ti2Ni and TiB is neglected. Turner’s model considered the internal stress system in a composite and calculated the CTE of the composite base on the bulk modulus of whisker and matrix. In Kerner’s model, it is assumed that the reinforcement particles were spherical and dispersed in the matrix. The CTE of the composite was calculated by taking both the normal and shear stresses into account. The mechanical constants (K, G, and α) of Ti2Ni matrix and TiB whiskers are listed in Table 4. Fig. 12 shows the CTE curves of the composite by different theoretical models. The Vw of TiB whiskers in the brazing seam was processed and calculated using Image J software. The Vw of TiB whiskers in the brazing seam was roughly estimated as 0.4 when the joint was brazed at 1040 °C for 30 min. The CTE of the brazing seam could be calculated as 27.27 × 10−6 K−1 for the Rom’s model, 24.50 × 10−6 K-1 for Turner’s model, and 25.63 × 10−6 K−1 for the Kerner’s model. The results indicated that the TiB whiskers could lower the CTE of the brazing seam. The CTE mismatch between the brazing seam and TTS ceramics could be reduced to relieve the residual stress of the brazed joints. As a result, a high joint strength could be achieved when the joint was brazed at 1040 °C for 30 min. Fig. 13 shows the fracture paths and morphologies of TTS ceramic brazed joints after the shear tests at room temperature and high temperature of 800 °C. When the joints are brazed at 1000 °C, cracks propagated in the TTS ceramic, and a classical bowed crack path was observed, as shown in Fig. 13(a). The result indicated that large residual stress was generated in the TTS ceramic. A huge mismatch of CTE between TTS ceramic and continuous Ti2Ni phase would lead to large residual stress in the brazed joint. The reaction phases formed in the brazing seam had a better plastic deformation behaviour than TTS ceramic. The reaction phases can release part of the stress in the brazing seam by its deformation. Thus, a large stress concentration was generated in the TTS ceramics, which led to the failure of the joint in the TTS ceramic. In contrast, when the joints are brazed at 1040 °C, a large number of fine TiB whiskers are formed in the brazing seam, as shown

Fig. 12. CTE curves of the composite by different theoretical models as a function of TiB volume fraction.

in Fig. 8(c). The TiB whiskers could reduce the CTE mismatch between the brazing seam and TTS ceramic, and therefore relieve the residual stress of the brazed joint. Under this circumstance, the failure of the brazed joint occurred at the brazing seam, as shown in Fig. 13(c). On the other hand, if the cracks have occurred at the brazing seam, the presence of fine TiB whiskers in the brazing seam also resulted in the deflection of the crack growth path, thus increasing energy dissipation in crack propagation. The tendency for cracks to deflect through TiB whiskers suggested that the joint shear strength could be improved by increasing the number of TiB whiskers. Therefore, a high joint strength was obtained after a shear test when the joints were brazed at 1040 °C. It can be seen from Fig. 13(d) that a large number of Ti2Ni phases and TiB whiskers were found in the fracture surface; this also confirms the above analysis of the fracture. Fig. 13(e)–(f) show the typical fracture path and morphology of the joint tested at 800 °C. Similar to the joint tested at RT, the failure has occurred in the brazing seam, and fractured surface is occupied by Ti2Ni phases and TiB whiskers. Owing to the melting point of Ti2Ni at 985 °C, the Ti2Ni phase was the weakest part of the brazed joint during the high-temperature shear test. Thus, the shear strength decreased sharply when the brazed joint was tested at elevated temperatures.

Table 4 Parameters of TiB whisker and Ti2Ni matrix used for the thermal expansion models calculation. Materials

K/GPa

G/GPa

α/10−6 K-1

TiB whisker [49,50] Ti2Ni matrix [45]

205.4 146.1

193.3 49.4

7.45 40.48

4. Conclusions TiB2-TiC-SiC ceramics were fabricated and joined using the brazing technique, which included the use of Ti-Ni composite foils. Further, the microstructure and mechanical properties of TiB2-TiC-SiC ceramics 9

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Fig. 13. Fracture paths and morphologies of the TTS ceramics brazed joints: (a–b) 1000 °C for 30 min after the RT shear test, (c–d) 1040 °C for 30 min after RT shear test, and (e–f) 1040 °C for 30 min after 800 °C shear test.

brazed joints were investigated. Based on the results obtained in this study, the main conclusions have been listed as follows:

ceramics and Ti became vigorous. The thickness of the reaction layer I increased with increasing the brazing temperature. In addition, the continuous Ti2Ni layer in the brazing seam gradually disappeared, and it was replaced by TiB and Ti5Si3. The maximum RT shear strength reached 168 MPa when the joint was brazed at 1040 °C for 30 min. Conversely, it was 104 and 81 MPa at test temperatures of 600 °C and 800 °C, respectively. (4) The fracture occurred primarily at the TiB2-TiC-SiC ceramics when the joint was brazed at a low temperature. However, the fracture path changed to the brazing seam when the brazing temperature was 1040 °C. The change of the fracture path was primarily attributed to an increase in the content of TiB whiskers in the brazing seam. The TiB whiskers could lower the CTE of the brazing seam and reduce a CTE mismatch between the brazing seam and the TiB2-TiC-SiC ceramic; this was probably relieved the residual stress of the brazed joint.

(1) TiB2-TiC-SiC ceramics primarily consisted of TiB2 matrix, off-white TiC, and dark grey SiC; it was observed that the TiB2, TiC, and SiC were uniformly distributed. The density of the hot-pressed TTS ceramic was 4.26 g/cm3, corresponding to about 98%. The flexural strength of the TiB2-TiC-SiC ceramic was 664 ± 32 MPa. (2) The typical interfacial structure of the TiB2-TiC-SiC ceramics brazed joint was primarily composed of a reaction layer I and brazing seam. Ti in the filler reacted with the TiB2-TiC-SiC ceramics and formed a reaction layer consisting of TiB and TiC, and the brazing seam was composed of TiB, TiC, Ti5Si3, Ti2Ni, and TiNi. (3) With an increase in the brazing temperature, the activity of Ti in the filler was gradually improved; the reaction between TiB2-TiC-SiC 10

Journal of the European Ceramic Society xxx (xxxx) xxx–xxx

X.Q. Cai, et al.

Declaration of Competing Interest

[22] L. Esposito, D. Sciti, L. Silvestroni, C. Melandri, S. Guicciardi, N. Saito, K. Nakashima, A.M. Glaeser, Transient liquid phase bonding of HfC-based ceramics, J. Mater. Sci. 49 (2014) 654–664. [23] N. Saito, H. Ikeda, Y. Yamaoka, A.M. Glaeser, K. Nakashima, Wettability and transient liquid phase bonding of hafnium diboride composite with Ni-Nb alloys, J. Mater. Sci. 47 (2012) 8454–8463. [24] W. Yang, T. Lin, P. He, H. Wei, L. Xing, D. Jia, Microstructure and mechanical properties of ZrB2-SiC joints fabricated by a contact-reactive brazing technique with Ti and Ni interlayers, Ceram. Int. 40 (5) (2014) 7253–7260. [25] W. Yang, L. Xing, T. Lin, P. He, J. Lin, X. Ma, Microstructural evolution and growth/ degradation behavior of in situ TiB whiskers in ZrB2-SiC joints using Ti/Ni/Ti filler, J. Alloys Compd. 744 (2018) 124–131. [26] R. Asthana, M. Singh, Joining of ZrB2-based ultra-high-temperature ceramic composites using Pd-based braze alloys, Scr. Mater. 61 (2009) 257–260. [27] F.C. Wang, Z.H. Zhang, L. Jie, C.C. Huang, S.-K. Lee, A novel rapid route for in situ synthesizing TiB-TiB2 composites, Compos. Sci. Technol. 69 (2009) 2682–2687. [28] Y.F. Zhang, Z.W. Li, C.G. Li, Z.H. Yu, Temperature gradient field sintering of Ti-TiBTiB2 functionally graded material, Ceram. Int. 41 (2015) 13844–13849. [29] J.M. Shi, J.C. Feng, X.Y. Tian, H. Liu, L.X. Zhang, Interfacial microstructure and mechanical property of ZrC-SiC ceramic and Ti6Al4V joint brazed with AgCuTi alloy, J. Eur. Ceram. Soc. 37 (8) (2017) 2769–2778. [30] J.M. Shi, L.X. Zhang, X.Y. Pan, X.Y. Tian, J.C. Feng, Microstructure evolution and mechanical property of ZrC-SiC/Ti6Al4V joints brazed using Ti-15Cu-15Ni filler, J. Eur. Ceram. Soc. 38 (4) (2018) 1237–1245. [31] Z.R. Li, Z.Z. Wang, G.D. Wu, J.C. Feng, Microstructure and mechanical properties of ZrB2-SiC ultrahigh temperature ceramic composite joint using TiZrNiCu filler metal, Sci. Technol. Weld. Joi. 16 (8) (2013) 697–701. [32] G. Wang, P. Xiao, Z. Huang, R. He, Brazing of ZrB2-SiC ceramic with amorphous CuTiNiZr filler, Ceram. Int. 42 (4) (2016) 5130–5135. [33] R.K. Shiue, S.K. Wu, S.Y. Chen, Infrared brazing of TiAl intermetallic using BAg-8 braze alloy, Acta Mater. 51 (2003) 1991–2004. [34] X.Y. Tian, J.C. Feng, J.M. Shi, Y.Z. Liu, L.X. Zhang, Interfacial microstructure and mechanical properties of the vacuum brazed C/SiC composite and Nb joints, Vacuum 146 (2017) 97–105. [35] Z. Chen, M.S. Cao, Q.Z. Zhao, J.S. Zou, Interfacial microstructure and strength of partial transient liquid-phase bonding of silicon nitride with Ti/Ni multi-interlayer, Mater. Sci. Eng. A 380 (2004) 394–401. [36] W.C. Oliver, G.M. Pharr, Measurement of hardness and elastic modulus by instrumented indentation: advances in understanding and refinements to methodology, J. Mater. Res. 19 (1) (2011) 3–20. [37] C.R. Weinberger, G.B. Thompson, Review of phase stability in the group IVB and VB ransition-metal carbides, J. Am. Ceram. Soc. 101 (2018) 4401–4424. [38] B.G. Fu, H.W. Wang, C.M. Zou, Z.J. Wei, Microstructural characterization of in situ synthesized TiB in cast Ti-1100-0.10B alloy, Trans. Nonferrous Met. Soc. China 25 (2015) 2206–2213. [39] W.J. Lua, D. Zhang, X.N. Zhang, R.J. Wu, T. Sakata, H. Mori, Microstructural characterization of TiB in in situ synthesized titanium matrix composites prepared by common casting technique, J. Alloys Compd. 327 (2001) 240–247. [40] H.B. Feng, Y. Zhou, D.C. Jia, Q.C. Meng, J.C. Rao, Growth mechanism of in situ TiB whiskers in spark plasma sintered TiB/Ti metal matrix composites, Cryst. Growth Des. 6 (7) (2006) 1626–1630. [41] Q. Meng, H. Feng, G. Chen, R. Yu, D. Jia, Y. Zhou, Defects formation of the in situ reaction synthesized TiB whiskers, J. Cryst. Growth 311 (6) (2009) 1612–1615. [42] Y. Song, D. Liu, S. Hu, X. Song, J. Cao, Graphene nanoplatelets reinforced AgCuTi composite filler for brazing SiC ceramic, J. Eur. Ceram. Soc. 39 (4) (2019) 696–704. [43] M. Singh, T. Matsunaga, H.-T. Lin, R. Asthana, T. Ishikawa, Microstructure and mechanical properties of joints in sintered SiC fiber-bonded ceramics brazed with Ag-Cu-Ti alloy, Mater. Sci. Eng. A 557 (2012) 69–76. [44] Y.L. Wang, W.L. Wang, J.H. Huang, R.H. Yu, J. Yang, S.H. Chen, Reactive composite brazing of C/C composite and GH3044 with Ag-Ti mixed powder filler material, Mater. Sci. Eng. A 759 (2019) 303–312. [45] D. Toprek, J. Belosevic-Cavor, V. Koteski, Ab initio studies of the structural, elastic, electronic and thermal properties of NiTi2 intermetallic, J. Phys. Chem. Solids 85 (2015) 197–205. [46] H.T. Vo, M. Todd, F.G. Shia, A.A. Shapiro, M. Edwards, Towards model-based engineering of under®ll materials: CTE modeling, Microelectronics J. 32 (2001) 331–338. [47] W. Zheng, X. He, M. Wu, S. Ren, S. Cao, D. Guan, R. Liu, X. Qu, Thermal expansion coefficient of Diamond/SiC composites prepared by silicon vapor infiltration in vacuum, Vacuum 159 (2019) 507–515. [48] E.H. Kerner, The elastic and thermo-elastic properties of composite media, Proc. Phys. Soc. Lond. B 69 (8) (1956) 808. [49] L. Sun, Y. Gao, B. Xiao, Y. Li, G. Wang, Anisotropic elastic and thermal properties of titanium borides by first-principles calculations, J. Alloys Compd. 579 (2013) 457–467. [50] K.S. Ravi Chandran, K.B. Panda, S.S. Sahay, TiBw-reinforced Ti composites: processing, properties, application prospects, and research needs, JOM 5 (2004) 42–48.

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements The authors gratefully acknowledge the financial support from the National Natural Science Foundation of China (Grant No. 51774214) and Key Laboratory of Advanced Ceramics and Machining Technology, Ministry of Education (Tianjin University). References [1] W.G. Fahrenholtz, G.E. Hilmas, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (5) (2007) 1347–1364. [2] S.R. Levine, E.J. Opila, M.C. Halbig, J.D. Kiser, M. Singh, J.A. Salem, Evaluation of ultra-high temperature ceramics for aeropropulsion use, J. Eur. Ceram. Soc. 22 (2002) 2757–2767. [3] W.G. Fahrenholtz, G.E. Hilmas, Ultra-high temperature ceramics: materials for extreme environments, Scr. Mater. 129 (2017) 94–99. [4] R. Savino, L. Criscuolo, G.D. Martino, S. Mungiguerra, Aero-thermo-chemical characterization of ultra-high-temperature ceramics for aerospace applications, J. Eur. Ceram. Soc. 38 (2018) 2937–2953. [5] A. Bhattacharya, C.M. Parish, T. Koyanagi, C.M. Petrie, D. King, G. Hilmas, W.G. Fahrenholtz, S.J. Zinkle, Y. Katoh, Nano-scale microstructure damage by neutron irradiations in a novel Boron-11 enriched TiB2 ultra-high temperature ceramic, Acta Mater. 165 (2019) 26–39. [6] E. Sani, M. Meucci, L. Mercatelli, A. Balbo, C. Musa, R. Licheri, R. Orrù, G. Cao, Titanium diboride ceramics for solar thermal absorbers, Sol. Energy Mater. Sol. Cells 169 (2017) 313–319. [7] S. Nekahi, M. Vajdi, F.S. Moghanlou, K. Vaferi, A. Motallebzadeh, M. Özen, Umut Aydemir, Jianjun Sha, Mehdi Shahedi Asl, TiB2-SiC-based ceramics as alternative efficient micro heat exchangers, Ceram. Int. 45 (2019) 19060–19067. [8] R. Königshofer, S. Fürnsinn, P. Steinkellner, W. Lengauer, R. Haas, K. Rabitsch, M. Scheerer, Solid-state properties of hot-pressed TiB2 ceramics, Int. J. Refract. Met. Hard Mater. 23 (4-6) (2005) 350–357. [9] A. Mukhopadhyay, T. Venkateswaran, B. Basu, Spark plasma sintering may lead to phase instability and inferior mechanical properties: a case study with TiB2, Scr. Mater. 69 (2) (2013) 159–164. [10] A. Tampieri, A. Bellosi, Oxidation of monolithic TiB2 and of Al2O3-TiB2 composite, J. Mater. Sci. 28 (3) (1993) 649–653. [11] B. Zou, W. Ji, C. Huang, K. Xu, S. Li, Degradation of strength properties and its fracture behaviour of TiB2-TiC-based composite ceramic cutting tool materials at the high temperature, Int. J. Refract. Met. Hard Mater. 47 (2014) 1–11. [12] A. Mukhopadhyay, G.B. Raju, B. Basu, A.K. Suri, Correlation between phase evolution, mechanical properties and instrumented indentation response of TiB2-based ceramics, J. Eur. Ceram. Soc. 29 (3) (2009) 505–516. [13] G.L. Zhao, C.Z. Huang, H.L. Liu, B. Zou, H.T. Zhu, J. Wang, Microstructure and mechanical properties of hot pressed TiB2-SiC composite ceramic tool materials at room and elevated temperatures, Mater. Sci. Eng. A 606 (2014) 108–116. [14] G.B. Raju, B. Basu, Densification, sintering reactions, and properties of titanium diboride with titanium disilicide as a sintering aid, J. Am. Ceram. Soc. 90 (11) (2007) 3415–3423. [15] D. Vallauri, I.C. Atías Adrián, A. Chrysanthou, TiC-TiB2 composites: a review of phase relationships, processing and properties, J. Eur. Ceram. Soc. 28 (8) (2008) 1697–1713. [16] D.S. King, W.G. Fahrenholtz, G.E. Hilmas, Silicon carbide–titanium diboride ceramic composites, J. Eur. Ceram. Soc. 33 (15–16) (2013) 2943–2951. [17] F. Mestral, F. Thevenot, Ceramic composites: TiB2-TiC-SiC, J. Mater. Sci. 26 (1991) 5561–5565. [18] G. Zhao, C. Huang, N. He, H. Liu, B. Zou, Microstructure and mechanical properties at room and elevated temperatures of reactively hot pressed TiB2-TiC-SiC composite ceramic tool materials, Ceram. Int. 42 (4) (2016) 5353–5361. [19] D.S. King, G.E. Hilmas, W.G. Fahrenholtz, J. Garay, Plasma arc welding of TiB2-20 vol% TiC, J. Am. Ceram. Soc. 97 (1) (2014) 56–59. [20] W. Yang, T. Lin, P. He, M. Zhu, C. Song, D. Jia, J. Feng, B. Derby, Microstructural evolution and growth behavior of in situ TiB whisker array in ZrB2-SiC/Ti6Al4V brazing joints, J. Am. Ceram. Soc. 96 (12) (2013) 3712–3719. [21] X.Y. Tian, J.C. Feng, J.M. Shi, H.W. Li, L.X. Zhang, Brazing of ZrB2-SiC-C ceramic and GH99 superalloy to form reticular seam with low residual stress, Ceram. Int. 41 (2015) 145–153.

11