Microstructural characterization and mechanical properties of a novel TiC-based cermet bonded with Ni3(Al,Ti) and NiAl duplexalloy

Materials Characterization 135 (2018) 295–302

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Microstructural characterization and mechanical properties of a novel TiCbased cermet bonded with Ni3(Al,Ti) and NiAl duplexalloy

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Wenqiang Hua, Zhenying Huanga,b, , Leping Caia, Cong Leia, Wenbo Yua, Hongxiang Zhaia, Shaoshuai Woa, Xinkang Lia, Yang Zhoua Institute of Materials Science and Engineering, School of Mechanical and Electronic Control Engineering, Beijing Jiaotong University, Beijing 100044, China Key Laboratory of Vehicle Advanced Manufacturing, Measuring and Control Technology (Beijing Jiaotong University), Ministry of Education, Beijing 100044, China

A R T I C L E I N F O

A B S T R A C T

Keywords: TiC-based cermets Ti3AlC2 Microstructure Ni3(Al,Ti) NiAl High-temperature properties

A novel TiC/Ni3(Al,Ti)-NiAl composite has been designed and synthesized by in-situ reactive hot-pressing a mixture of Ti3AlC2 and Ni powders. After hot-press sintering at 1500 °C and 30 MPa, almost fully dense composite was obtained and the in-situ formed TiC grains were well-bonded with Ni3(Al,Ti) and NiAl alloys. Due to the nanolaminated structure in Ti3AlC2 materials that weakly bonded Al layers are more prone to be out-diffused from Ti3AlC2 lattice than that of Ti atoms. The deintercalated AleTi atoms would react with Ni to give rise to the formation of Ni3(Al,Ti) and NiAl alloys, and meanwhile a crystal lattice structure transformation happened from hexagonal Ti3AlC2 to cubic lattice TiC. The formation of phases, grains morphologies, sintering behavior and mechanical properties were associated with sintering temperature. The as prepared TiC/Ni3(Al,Ti)-NiAl composite under 1500 °C exhibited superior mechanical properties comparing with the counterparts fabricated under 1350 °C and 1400 °C because of the enhanced microstructures and eliminations of pores. Hardness, flexural strength and fracture toughness of the TiC/Ni3(Al,Ti)-NiAl composite at room temperature were determined as 9.9 ± 0.35 GPa, 665 ± 26 MPa and 10.23 ± 0.4 MPa·m1/2, respectively. Interestingly, the flexural strength and fracture toughness increases with temperature rising from 600 to 800 °C, the maximum of flexural strength and fracture toughness could reach 775 ± 25 MPa and 11.6 ± 0.4 MPa·m1/2 at 800 °C.

1. Introduction Much research efforts during recent years have been devoted into the field of structural materials of composite with the promise of multifunctionality or a distinct spectrum of specific properties and possibility of realizing unique combination of performances that is unachievable with conventional and simplex materials [1,2]. Among those advanced materials, the hard brittle materials like ceramic bonded with soft metal or intermetallic phase (also called “cermet”) with fine wettability were developed to be as the promising candidate used in high-temperature service area [3–5]. In the last few years, tungsten carbide‑cobalt (WC-Co) was successfully applied as the cutting tools [6]. Nevertheless, these materials exhibited unsatisfied oxidation and corrosion resistance, low susceptibility to plastic deformation at high temperature as well as sparse and expensive Co material, which consequently limited their further practical applications [7–9]. For the purpose of improving corrosion and oxidation resistance and replace costly Co, titanium carbide based cermets with a great number of ductile metals and intermetallics like Ni3Al and Ni2AlTi as the binding

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materials have been extensively investigated [10–19]. They are known to possess useful properties such as relatively lower density and better elevated temperature strength, exceptional corrosion resistance in oxygen and carbon enriched atmosphere and resistance in organic acids, bases and sodium-chloride solution environment. β-NiAl exhibits low density, high melting point, good high-temperature oxidation and corrosion resistance. However, the lack of room temperature ductility and low elevated-temperature strength have restricted its practical applications [20]. In the past few years, many investigations revealed that microstructure modifications of β-NiAl including the ductile γ′ phase (Ni3Al or Ni3(Al,Ti)) resulted in significant improvement of hightemperature strength and room temperature ductility [21–24]. To obtain γ′ phase in the NiAl matrix, Ti3AlC2 and Ni characterized as precursors were adopted in the present study, resulting in the formation of TiC-based cermets bonded with γ′ + β duplexalloy. Titanium aluminum carbide (Ti3AlC2), as one of the most attractive ternary carbides in the Mn + 1AXn (n = 1, 2 or 3) phase family (M is an early transition metal, A is an A-group element, and X is carbon and/or nitrogen), has attracted huge attention in virtue of its intriguing

Corresponding author at: Institute of Materials Science and Engineering, School of Mechanical and Electronic Control Engineering, Beijing Jiaotong University, Beijing 100044, China. E-mail address: [email protected] (Z. Huang).

https://doi.org/10.1016/j.matchar.2017.11.003 Received 26 September 2017; Received in revised form 2 November 2017; Accepted 2 November 2017 Available online 09 November 2017 1044-5803/ © 2017 Elsevier Inc. All rights reserved.

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Fig. 1. SEM images of starting materials: (a) Electrolytic Ni powders; (b) Homemade Ti3AlC2 powders.

Fig. 2. XRD patterns of Ti3AC2-Ni reactant powders after sintering at 1350 °C, 1400 °C and 1500 °C for 30 min at Ar atmosphere, respectively.

Fig. 3. The standard Gibbs free energies of formation for Ni3(Al,Ti), NiAl, TiC and NiTi.

The de-intercalated Al and fractional Ti atoms would primarily react with Ni to form Ni3(Al, Ti) (L12 structure) due to the lower Gibbs free energy of γ′ phase and revolve around the topotactic formation of TiC grains. Subsequently, due to high diffusivity of Al atoms, the remanent Al atoms will keep on diffusing outward to react with external Ni to give rise to the formation of NiAl (B2 structure). Hence, a novel TiCbased cermet bonded with γ′ + β duplexalloy that potentially has lower density and better oxidation resistance than the metal (Ni or/and Co), better fracture and damage resistance than the ceramic (TiC) was in-situ fabricated with starting materials of Ti3AlC2 and Ni powders. However, there is no report on the TiC/Ni3(Al,Ti)-NiAl composite so far. Studies of the elevated-temperature strength of carbide composites based on TiC, in particular those bonded with intermetallics appear to be infrequent. In conclusion, in the present study, a new family of TiC-based cermet have been prepared with a nominally stoichiometric Ni3(Al,Ti) and NiAl duplexally binders through a novel in-situ reaction between Ti3AlC2 and Ni starting materials during sintering process. The effect of sintering temperature (1350 °C, 1400 °C, 1500 °C) on microstructure and mechanical properties of samples has been examined, paying particular attention to the sample sintered at 1500 °C, which presents almost a fully dense and enhanced microstructure. Some performance characteristics - Vicker's hardness, flexural strength and fracture toughness at RT and elevated temperature (from 200 °C to 1000 °C) were determined. The reaction mechanism in the Ti3AlC2-Ni system was explicated correspondingly together with standard Gibbs free energies (Gibbs) of each formation of products and a reaction model to better comprehend the reaction process was proposed and finally interpreted the fracture mechanism of TiC/Ni3(Al,Ti)-NiAl composite.

combination of ceramic and metallic properties at both room and elevated temperatures [25–29]. Similar to the vast majority of the MAX phases, Ti3AlC2 (hexagonal with a = 3.0753 Å and c = 18.578 Å) is excellent conductor of electricity and heat and most readily machinable, with low susceptibility to thermal shock, quite stiff and lightweight and exceptionally damage tolerant [30–32]. Moreover, Ti3AlC2 materials are distinguished by a unique combination of the properties which are derived from their nanolayered structure of strong covalent bonded TieC layers interleaved with Al metal layers through weaker TieAl bonds [33–35]. Herein, the instability of the Ti3AlC2 crystal lattice tends to liberate Al atoms and leave vacancies at Al sites behind. This unique characteristic of Ti3AlC2 and other MAX phase materials have also been demonstrated by many investigations [36–41]. For example, Naguib et al. [37,38] discovered that by immersing Al-containing Ti3AlC2 in HF acid, it was possible to selectively etch the Al, resulting in the formation of exfoliated two-dimensional, 2D, Ti3C2 layers labeled as MXene, which, in essence, makes the connection to another conducting 2D graphene material. Wang et al. [39] demonstrated that when Ti3AlC2 bulks are heated up to 1100 °C in air, the surface of sample is covered with a layer of Al2O3, meanwhile, a crystal lattice structure transformation happened from hexagonal to cubic lattice, a part of Ti3AlC2 substrate decomposes into TiCx because of deintercalation of enough Al from Ti3AlC2. Zhang et al. [19] reported a novel twining platelets strengthened TiC-Ni2AlTi composite fabricated by in situ reactive hot-pressing method of blended Ti3AlC2 and Ni powders. Based on such speciality of MAX phase materials that in-situ TiCx reinforced MMCs (like TiC0.61/Cu(Al) [42], TiCx/Fe(Al) [43], TiC0.53/Ni(Si,Ti) [44], TiCx-Ni3(Al,Ti)/Ni [45] etc.) with enhanced properties have also been widely developed by using Ti3AlC2 or Ti3SiC2 as precursors. Particularly, according to our previous investigations in Ti3AlC2-Ni system, the molten Ni could trigger the decomposition of Ti3AlC2[45]. 296

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Fig. 4. Microstructure of as-prepared composites sintered at 1350 °C, 1400 °C and 1500 °C, respectively: (a) and (b) 1350 °C sintered; (c) and (d) 1400 °C sintered; (e) and (f) 1500 °C sintered.

2. Experimental Procedures

1000 °C, as well as Vicker's hardness. All the specimens for properties determination were machined by wire electrode cutting, and then polished to minimize the effects of surface irregularities and finish, using emery papers impregnated with silicon carbide from 240 to 2000 grit, and finally cleaned thoroughly using ultrasonic bath and ethanol. Specimens for flexural strength examinations were machined into bend bar with nominal dimensions of 3 mm × 4 mm × 36 mm and determined in three-point bending test with span of 30 mm on a universal electron material testing machine (WDW-100E, Shidai, China) under a loading speed of 0.5 mm/min at room temperature. The flexural strength at elevated temperature was measured at 200 °C, 400 °C, 600 °C, 800 °C, 900 °C and 1000 °C, respectively, in air after soaking temperature for 10 min using universal electron material testing machine equipped with high temperature furnace (GW-1200, Fangrui, China) at a heating rate of 10 °C/min. The fracture toughness (K1C) of the sintered composites was measured by edge-notched bending tests. The testing specimens with dimension of 2.5 mm × 5 mm × 26 mm with a notch of 0.2 mm in width and 2.5 mm in depth were used for the evaluation of fracture toughness by universal electron material testing machine under a loading speed of 0.05 mm/min. In addition, the measurement of fracture toughness under high temperature was similar with the flexural strength aforementioned but for the load speed, which is 0.05 mm/min for K1C. The hardness was determined using Vicker's hardness tester (TH-700, Shidai, China) in a load range of 1–10 kg at a dwelling contact time of 15 s. The density of as prepared samples was measured by Archimedes method based on the ASTM C20-00.

The starting materials for investigations of TiC-based cermet were homemade Ti3AlC2 (purity 98%, 6–8 μm in length and 3–5 μm in width) [46] and commercial electrolytic Ni powders (purity 99.3%, particle size < 50 μm, General Research Institute for Nonferrous Metals, Beijing, China). The morphologies of raw material powders are shown in Fig.1. The enlarged image of rectangular area marked in Fig.1(b) illustrates the emblematical laminated structure of ternary Ti3AlC2. The mixtures of Ti3AlC2 and Ni were ball-milled for 10 h before sintering with a ball-to-powder weight ratio of 10:1 at a speed of 300 rpm. The milling was performed in a full-ranged planetary ball mill (QM-3SP4, Nanjing university instrument plant, China). Stainless steel ball and jar were adopted in the process of mechanical milling. The homogenized powders were monaxially compressed into a green compact in a graphite die coated with h-BN. Afterwards, the green compact together with the mold were placed into vacuum hot-pressing furnace (ZT-30-22Y, Chenhua, China) and heated to 1350 °C, 1400 °C and 1500 °C with soaking time of 30 min, respectively, at a heating rate of 10 °C/min in argon atmosphere. Phase evolutions of the sintered samples were identified by X-ray diffractometer (XRD) using Cu Kα radiation (X′ PERT-PRO MPDT, Netherlands). The microstructure and phase compositions of the composites were assessed using a ZEISS EVO 18 scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer system (EDS). The Gibbs free energies of formed products were calculated by the HSC Chemistry 6.0 software. Mechanical properties of synthesized composite were evaluated in term of flexural strength and fracture toughness determined from RT to 297

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phase is ascertained. No residual trace of Ti3AlC2 phase can be identified, indicating that the decomposition of Ti3AlC2 has been entirely completed. As indicated, TiC, Ni3(Al,Ti), NiAl and NiTi are detected as dominating reaction products after sintering at 1350 °C. However, when the sintering temperature increases to 1400 °C, the diffraction peaks of NiTi are attenuated obviously but still distinguishable and vanish finally after sintering at 1500 °C, revealing the metastable NiTi alloy has been decomposed and transformed into some stable phases. Thus, the final composite should mainly contain the hard TiC grains with the binding materials of accrete NiAl and Ni3(Al,Ti) alloys. The inset offset shows the evolution process of NiTi with temperature rising, which suggested that a strong temperature dependence of the stability of this metastable phase. The transitional phase like NieTi compound existing in the final products has also been discovered in TieCeNi systems through combustion synthesis [48]. Zhang et al. suggested that transitional phases like NiTi and Ni3Ti compounds would retain at a relatively low combustion temperature due to the considerable reduction of contact surface area between Ni and Ti, but only the TiC thermodynamically stable phases and Ni as a diluent exist when the reaction temperature was beyond the Ti melting point of 1672 °C [49]. NieTi was readily participated in the interfacial reactions in the TieNieC system, and the NieTi compounds were usually detectable in the interface of TiC grains [50]. In addition, the shoulder peaks at TiC are visible (the signature peaks are 2θ = 41.927 and 60.793, respectively) indicating the sub-stoichiometric TiCx formed. The occurrence can be attributed to the complete disintegration of NiTi alloy, which provided additional Ti element for TiC formation, resulted in the increase of Ti/C ratio, therefore, additional TiCx appearing at the shoulders of standard TiC peaks. The standard Gibbs free energies (△G) of products in the TieAleCeNi system were calculated by the HSC6.0 Chemistry software to verify the possibility of formation of products in the composite, see Fig. 3. The thermodynamic assessment suggests Ni3(Al,Ti) is the primary phase in the composite sintered from 1350 °C − 1500 °C due to the lower standard Gibbs free energy of Ni3(Al,Ti) formation (△Gf (1350 °C) = −492.28 kcal/mol) compared to that of TiC (△Gf (1350 °C) = −288.18 kcal/mol), NiAl (△Gf (1350 °C) = −277.98 kcal/mol) and NiTi (△Gf (1350 °C) = −231.82 kcal/ mol). The highest △Gf of NiTi formation in the products indicates it is a metastable phase or will not formed in this system. Whereas, the content of Ti in Ti3AlC2 is much higher than that of Al, the formation of initial NiTi is inevitable at the early stage of sintering, and diminishes finally with rising temperatures. In addition, it could be noted that theoretical Ti3C2 (C/Ti = 0.67)

Fig. 5. EDS results for the area in Fig. 4 (b). Panels (a) and (b) are corresponding to the results labeled as point “1” and “2”, respectively.

3. Results and Discussion 3.1. Phase Analysis The thermodynamic assessment of 900 °C isothemal section in the NieAleTi ternary phase diagram system revealed that Ni3(Al,Ti) (L12 structure) and NiAl (B2 structure) two phases should form based on the element composition and remained stable in a large area due to their relatively lower standard Gibbs free energy of formation [47]. Fig.2 illustrates the XRD patterns of Ti3AlC2 - Ni reactant mixtures after sintering at different temperatures. It should be observed that the three XRD patterns are very similar with each other, and TiC as the main

Fig. 6. BSE image of the polished surface with EDS line scans for the profile of different elements obtained after fabricated at 1500 °C. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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Fig. 7. A schematic diagram illustrating AleTi atoms out-diffusion from Ti3AlC2 and corresponding crystal lattice transformation of Ti3AlC2 from hexagonal to cubic lattice TiC.

Table 1 Mechanical properties of as-prepared samples made from Ti3AlC2 and Ni as starting materials. Samples

Density (g/cm3)

Relative density (%)

Hardness (30 N, GPa)

Flexural strength (MPa)

Fracture toughness (MPa·m1/2)

1350 °C 1400 °C 1500 °C

6.11 ± 0.19 6.05 ± 0.28 5.98 ± 0.12

91 ± 0.6 94 ± 0.75 97.4 ± 0.3

8.4 ± 0.48 8.69 ± 0.37 9.9 ± 0.35

578 ± 65 594 ± 42 665 ± 26

6.39 ± 0.395 7.56 ± 0.33 10.23 ± 0.38

dark grey particles embedded in a light grey matrix. Fig.4(a) and (b) are SEM images of polished surface of the sample sintered at 1350 °C. The sample mainly contains three kinds of different-contrast phases and some pores could also be observed. It can be clearly noted from the image that the well distributed Ti3AlC2 particles were transformed into angular TiC particles. In addition, other phases in back-scatted mode were determined according to the EDS results based on the atom ratios of the different elements and XRD analysis. In dark-grey area, the composition of elements are about 57 at.% Ni, 23 at.% Ti and 20 at.% Al, and light-grey color area contains mainly of Ni and Al elements, see Fig.5(a) and (b). Therefore, we can deduce that it is a two-phase region of Ni3(Al,Ti) and NiTi in dark grey area, and NiAl corresponding in light-grey color area. Besides, there are many pores existing in the cermet, indicating that sintering was limited and incomplete. Lower temperature brings about the imperfect penetration of binder phase to TiC particles. After that, at 1400 °C sintering shown in Fig.4(c), the amount of pores decreased comparatively. Fig.4(d) is the enlarged image of rectangular frame in Fig.4(c). The sintered composite had some TiC grains well-surrounded by a rim of Ni3(Al,Ti), forming a diverse core-rim structures. It is deserved to be mentioned that the corerim microstructure presented in this work is different from the general TiC-based cermet bonded with Ni or/and Co which a fair assessment of dissolution-reprecipitation formation mechanism was commonly approbatory [10–12]. In this work, each of Ti3AlC2 unit cell is a nucleating core of TiC, and the formation of rim characterized as Ni3(Al,Ti) is attributed to the de-intercalation of AleTi atoms from Ti3AlC2

should be obtained when Ti3AlC2 transformed topotactically. This topotactic transformation of Ti3AlC2 has also been confirmed in the previous studies that in-situ TiCx (x ≤ 0.7) was obtained to reinforce MMCs (Cu, Fe, Ni) [42–45]. Al atoms diffused into the liquid metal and Ti3AlC2 correspondingly dissociated into substoichiometric TiCx. Comparably, the stoichiometric TiC was obtained in the present study, the interpretation about this appearance may be ascribed to the primary formation of Ni3(Al,Ti). In addition, the sintering temperature plays a crucial role in the determination of the lattice parameters and stoichiometry of TiCx, the higher temperature thermodynamically prompting x values increasing [51]. 3.2. Microstructure Investigations The microstructural morphologies of the composites developed during sintering process depends dominantly upon the wetting behavior between alloys and ceramic particles. The higher sintered temperature (from 1350 °C to 1500 °C in this work) gives rise to densification occurring by essential liquid phase sintering procedure with particle rearrangement. In the final microstructure of synthesized TiC/ Ni3(Al,Ti)-NiAl composite, the NiAl alloy forms a semicontinuous intergranular binder with some remnant Ni3(Al,Ti)-rich areas revolving around the in-situ formation of TiC grains. Representative SEM microstructures evolvement of samples sintered under 1350 °C, 1400 °C and 1500 °C are shown in Fig.4(a)–(f). The overall microstructure of the composite appears rather uniform with 299

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atoms. It is the specificity of laminated structure in Ti3AlC2 materials that weakly bonded Al layers are more prone to be out-diffused from the Ti3AlC2 than that of Ti atoms, which covalently bonded Ti and C atoms in Ti6C octahedron retards the outward diffusion of Ti [25,27]. The de-intercalated Al and Ti atoms would react with Ni around the Ti3AlC2 precursor, prompting the formation of a rim of Ni3(Al,Ti). At this moment, Al is not exhausting yet, the remanent Al going on spread outside tends to react with Ni remaining to form NiAl alloy. The overall reaction procedure discussed above can be simplified by a model. Fig. 7 shows the diffusion process of AleTi atoms discussed above and simultaneously topotactic transformation of Ti3AlC2, in which a structural transformation from a hexagonal Ti3AlC2 to a cubic lattice of TiC with partial loss of layering can be clearly noted. However, it is well established that the ternary carbides and nitrides with a Mn + 1AXn formula form laminated structures with characteristic anisotropic properties [25]. A-site elements diffuse dominantly along the a-axis on the basal plane in the unit cell. If a-axis in Ti3AlC2 unit cell is parallel to the surface of sample, the AleTi atoms would spread evenly to the outside of Ti3AlC2 precursor and then react with Ni. Thus the well-defined Ni3(Al,Ti) rim can be observed distinctly, otherwise, partial and agglomerative Ni3(Al,Ti) accumulated at the interface of TiC grains (Seen from the Fig.7). Based on the results presented above it is reasonable to conclude that the following simplified reactions occur when Ti3AlC2 is reacted with Ni. Reaction (1) is essential and is followed by Reactions (2)–(4). In the previous studies, many researchers have proposed the evidence about the reaction mechanism between MAX phases (Ti3AlC2 and Ti3SiC2) and Ni [44,45]. In the present TieAleCeNi system, similar topotactic reaction (1) occurred to form in-situ TiC particles.

Ti3AlC2 = 2 Ti3 − x C2 (s ) + Al − x Ti (l)

(1)

Reactions (2) presented the formation of Ni3(Al,Ti) and NiTi during the sintering process. From the SEM images, it can be obviously observed that increasing number of TiC particles were generated with rising temperatures. These results can be interpreted by reaction (3), which NiTi disintegration happened and transformed into stable TiC and Ni melt [49]. Reaction (4) illustrates the formation of NiAl alloy.

Fig. 8. Mechanical properties TiC/Ni3(Al,Ti)-NiAl composite sintered under 1500 °C: (a) Vickers' hardness of of the composite as function of applied loads; (b) Temperature dependence of flexural strength and fracture toughness of TiC/Ni3(Al,Ti)-NiAl composites.

accompanied with following reaction between them and Ni. The details of these reaction procedures will be revealed in the afterwards EDS linescanning analysis. At 1500 °C, see from Fig.4(e) and (f), more densified and uniform microstructure, and no residual pores were observed. Higher temperature facilitates the liquid phase sintering which provided large contact areas between TiC and matrix phase (Ni3(Al,Ti) and NiAl) and promoted the process of diffusion of atoms along grain contacts and rapid mass transfer through the liquid, leading to grain shape accommodation and pores elimination. Fig. 6 shows the EDS line-scanning analysis of the sample obtained at 1500 °C, which the element distributions illustrating the diffusion process (labeled as red arrow) of AleTi atoms in the Ti3AlC2eNi system can be identified along the line. As indicated, the content of Al is increased gradually towards the direction of binder matrix while the content of Ti began to decline and remain stable at a low level. The interface between Ni3(Al,Ti) and NiAl can be distinguished by point “O” marked on the scanning line according to the distributions of AleTi

2 Al─Ti (l) + 4 Ni (l) = Ni3 (Al, Ti) (s ) + NiTi (s ) + Al (l)

(2)

NiTi (s ) + Ti3 − x C2 (s ) = TiC (s ) + Ni (l)

(3)

Ni (l) + Al (l) = NiAl (s )

(4)

From the above discussion, a final TiC/Ni3(Al,Ti)-NiAl composite was successfully synthesized from Ti3AlC2 and Ni starting powders via in-situ hot-pressing route at 1500 °C under 30 MPa. Then it is anticipated that these materials inherit excellent mechanical properties. 3.3. Mechanical Properties Room temperature mechanical properties of as prepared samples sintered under 1350 °C, 1400 °C and 1500 °C using Ti3AlC2 and Ni as starting materials were measured and presented in Table 1. The density Fig. 9. SEM images of indent on the TiC/Ni3(Al,Ti)NiAl composite prepared at 1500 °C and details of cracks propagation.

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Fig. 10. SEM image of the fractured TiC/Ni3(Al,Ti)NiAl composite fabricated at 1500 °C: (a) and (b) fracture morphology under room temperature; (c) and (d) fracture morphology under 800 °C.

With the combination of high hardness and excellent elevated temperature properties, the novel TiC/Ni3(Al,Ti)-NiAl composites reported in this work are possible condidates to be used as diesel engine components and cutting tools. Fig. 9 presents the Vicker's indents on the as-prepared sample generated with a load of 50 N. It is evident that the indentation is regular and without any crushing on the edge, instead, the clearly plastic deformation occurred. In addition, the cracks initiated at the corner of indentation. The details of the crack propagation are highlighted from the rectangular area in the indentation, in which the crack propagation is typified by an intergranular and transgranular mixed modes combined with deflection and bridging in the matrix. This multiple energyabsorbing mechanism will consume more strain energy and to some extent effectively explicating the determination of excellent fracture toughness. The surface morphologies of TiC/Ni3(Al,Ti)-NiAl composite after deformation at RT and 800 °C were examined by SEM. As indicated from Fig.10(a), fracture surface of the composite is mainly in the mode of transgranular form, but also accompanied by a few particles exhibiting an intergranular fracture. Fig. 10(b) shows an enlarged image of the area inside the rectangle in Fig. 10(a). Seen from Fig. 10(b), welldefined concavities corresponding to Ni3(Al,Ti) and some tear ridges can be clearly observed, suggesting the binder phase fails by interfacial decohesion between TiC and Ni3(Al,Ti). The formation of tearing ridges was due to the deformation of plastic Ni3(Al,Ti) and NiAl alloys. Besides, the failure of the matrix exhibits relatively flat facets that suggest cleavage as the additional failure mechanism. Compared with the fracture surface for cermet observed at ambient temperature, more rough surfaces were observed for the TiC/Ni3(Al,Ti)-NiAl composite determined at 800 °C, see Fig. 10(c). From the enlarged area in Fig. 10(c), Fig. 10(d) shows some transgranular microcracks though hard TiC grains. At high temperature, thermal stress at interface increases due to the mismatch of coefficient of thermal expansion between TiC and Ni3(Al,Ti), in which the interfacial tensile stress gives rise to the broken of TiC grains under the function of external load.

of samples tends to apparently increase with sintering temperatures increasing, and maximum density of 97.4% was observed in the case of specimen fabricated at 1500 °C. The effect of sintered temperature on sample densification depends mostly upon the amount of liquid phase formed in the TiC-based cermet system. Due to limited and incomplete liquid phase in the sintering process at 1350 °C and 1400 °C, the lower density values of 91% and 94% are eventually determined. Meanwhile, mechanical properties in term of hardness, flexural strength and fracture toughness are related to densification level of composites. With sintering temperature raising up, densification level of composites was enhanced resulting in the improvement of mechanical properties. Specifically, Hardness, flexural strength and fracture toughness improved from 8.4 ± 0.48 GPa, 578 ± 65 MPa and 6.39 ± 0.4 MPa·m1/2 to 9.9 ± 0.35 GPa, 665 ± 26 MPa and 10.23 ± 0.4 MPa·m1/2 with the sintering temperature rising from 1350 °C to 1500 °C. Fig. 8(a) shows the Vickers hardness of the TiC/Ni3(Al,Ti)-NiAl composite as a function of applied load. The Vickers hardness of the composite synthesized is about 10 GPa (at a load of 50 N) seen from the image. Fig. 8(b) plots the temperature dependence of flexural strength and fracture toughness of TiC/Ni3(Al,Ti)-NiAl composite. As indicated, taking the effect of erro bar into account, the strength and toughness values determined from RT to 400 °C can be considered similar (~ 665 MPa and 10.2 MPa·m1/2). When the testing temperature increased to 600 °C, the trends for flexural strength and fracture toughness tend to be opposite. After that, it is interesting to note that, in the investigated temperature range, the maximum strength and toughness values are obtained at 800 °C, which reach to 760 MPa and 11.5 MPa·m1/2, respectively. However, when the testing temperature exceeds 800 °C, strength and toughness decrease synchronously but more rapidly for flexural strength with the increase in temperature. The residual strength of the samples decreases from 650 MPa at room temperature to about 350 MPa as the testing temperature rises to 900 °C, and sequentially decreased to ~200 MPa at 1000 °C. In other words, the TiC/Ni3(Al,Ti)-NiAl composite could withstand until to 800 °C and without catastophic thermal failure yet when the testing temperature transcends this limitation. The reason for flexural strength of TiC/Ni3(Al,Ti)-NiAl composite increasing with rising temperature unusually which is similar to the yielding strength in the intermetallic compounds with L12 structure is likely attributed to Ni3(Al,Ti), which the strength, according to previous report [52,53], increased abnormally at the range of 600 °C–800 °C. However, for the present composite, interpretation of this phenomenon calls for further examinations.

4. Conclusion In this work, a novel TiC/Ni3(Al,Ti)-NiAl composite with the combination of excellent flexural strength, high fracture toughness and Vickers hardness was fabricated by reactive hot-press sintering of precursor powders of Ti3AlC2 and Ni under 1500 °C and 30 MPa. The 301

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molten Ni induced the decomposition of Ti3AlC2 under elevated temperature. The de-intercalated AleTi atoms would react with Ni to prompt the formation of Ni3(Al,Ti) and NiAl alloys, and meanwhile a crystal lattice structure transformation happened from hexagonal Ti3AlC2 to cubic lattice TiC. Hardness, flexural strength and fracture toughness of the composite at room temperature were determined as 9.9 ± 0.35 GPa, 665 ± 26 MPa and 10.23 ± 0.4 MPa·m1/2, respectively. The flexural strength and fracture toughness increases with temperature rising from 600 to 800 °C, the maximum of flexural strength and fracture toughness could reach 775 ± 25 MPa and 11.6 ± 0.4 MPa·m1/2 at 800 °C. The de-cohesion of TiC from Ni3(Al,Ti) interface and cleavage in the binder NiAl compound were determined as the dominating fracture mechanism. The excellent performance of TiC/Ni3(Al,Ti)-NiAl composite under high temperature also makes these materials as possible candidates for diesel engine components and cutting tools.

[20]

[21] [22] [23] [24]

[25] [26]

[27] [28]

Acknowledgments

[29]

This work was supported by National Science Foundation of China (NSFC) under Grant Nos. 51301013 and 51572017, by the Fundamental Research Funds for the Central Universities under Grant No. 2014JBZ015, and by Beijing Government Funds for the Constructive Project of Central Universities (353139535). The financial support by them are greatly appreciated.

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Appendix A. Supplementary data

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Supplementary data to this article can be found online at https:// doi.org/10.1016/j.matchar.2017.11.003.

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