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Characterization of intermetallic bonded TiC composites prepared by mechanically induced self-sustained reaction
Materials and Design 89 (2016) 102–108
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Materials and Design journal homepage: www.elsevier.com/locate/jmad
Characterization of intermetallic bonded TiC composites prepared by mechanically induced self-sustained reaction Xiao Chen a, Jianfeng Xu a,⁎, Quanquan Sun a, Wei Yang a, Weihao Xiong b a b
State Key Laboratory of Digital Manufacturing Equipment & Technology, School of Mechanical Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China
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Article history: Received 30 June 2015 Received in revised form 27 September 2015 Accepted 28 September 2015 Available online xxxx Keywords: TiC–Ni3Al Mechanically induced self-sustained reaction Microstructure Mechanical properties
a b s t r a c t The microstructure and mechanical properties of Ni3Al bonded TiC composites prepared by mechanically induced self-sustained reaction (MSR) were investigated. Ultrafine TiC–Ni3Al composite powders with various binder contents were synthesized by high energy ball milling of Ti–C–Ni–Al powder mixtures. Some porous blocks, which consisted of spherical TiC grains and Ni3Al binder, were also formed in each system after MSR. A core-rim structure present in the grains of the sintered TiC–Ni3Al composites was observed by Scanning Transmission Electron Microscopy (STEM). TiC particles dissolved selectively and then precipitated along with a stable crystal plane and gradually evolved into a faceted octahedron in the 3D space, which led to the formation of sharp-edged ceramic grains. High Resolution Transmission Electron Microscopy (HRTEM) analysis showed a 1–2 nm width transition zone and a low-angle misorientation between the ceramic phase and the binder phase. The TiC particles that precipitated from Ni3Al during MSR were easily wetted by the liquid Ni3Al binder, which gave rise to the superior mechanical properties of the TiC–Ni3Al composites. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Nickel aluminide (Ni3Al) has attracted considerable attention as an alternative binder for ceramic matrix composites due to its exceptional high temperature strength and excellent chemical stability [1–3]. The brittleness of Ni3Al appears to have been solved by doping boron [4– 6] and the good wettability of Ni3Al on TiC [7,8] makes it possible to use Ni3Al as a binder for TiC-based composites. To date, extensive attempts have been made to prepare TiC–Ni3Al composites [1,9–16]. This type of Ni3Al bonded TiC composites exhibit good oxidation and corrosion resistance, as well as high strength at elevated temperatures [1,16]. As such, they are promising materials for applications in high temperature, corrosion and wear environments, including cutting tools, metal heat forming tools and turbine blades for gas-turbine engines. However, the inferior flexure strength of TiC–Ni3Al composites at room temperature limits their practical application. Melt-infiltration sintering (MIS) has been adopted to prepare TiC– Ni3Al composites [1,9,11,15–17]. Some researchers have also prepared TiC–Ni3Al composites using vacuum sintering [7,10], reactive sintering [14,18] and self-propagating high temperature synthesis (SHS) [12, 13]. Mechanically induced self-sustaining reactions (MSR) have been used as a powerful manufacturing technique for the synthesis of a great number of novel materials [19–24]. This reactive milling method takes advantage of the strong exothermic character of the formation ⁎ Corresponding author. E-mail address: [email protected] (J. Xu).
http://dx.doi.org/10.1016/j.matdes.2015.09.158 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
of compounds from chemical elements to promote self-propagating reactions during milling and obtain particulate materials with high purity and homogenous composition. MSR has been shown to offer many advantages over other methods in terms of low cost, small particle sizes, and narrow size distribution products [20]. Moreover, the structure of these particles steadily diminishes and many dislocations are created during high-energy milling, which improves the mechanical properties of these novel materials [20,21]. In this reported work, various TiC– Ni3Al composites with binder contents from 20 wt.% to 50 wt.% were prepared using MSR with subsequent vacuum sintering. The synthesis of composite powders together with the microstructure and mechanical properties of the TiC–xNi3Al composites was investigated. 2. Experimental procedures Ti powder (99.0% pure, 45 μm), Ni powder (99.8% pure, 2.25 μm), Al powder (99.8% pure, 1.80 μm), graphite powder (99.9% pure, 5.50 μm)
Table 1 Target compositions, ignition time and lattice parameter of the synthesized Ni3Al. Mixture
Target composition (wt.%)
Ignition time
Lattice parameter (nm)
P1 P2 P3 P4 P5
TiC–20Ni3Al TiC–30Ni3Al TiC–40Ni3Al TiC–50Ni3Al Ni3Al
93 min 121 min 155 min 184 min –
0.3608 0.3602 0.3599 0.3594 0.3582
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Fig. 1. XRD patterns of the powders and blocks formed in Ti–C–Ni–Al system after MSR: (a) the powders formed in P1–P5; (b) the blocks formed in P1–P4.
and B powder (99.0% pure, 1–15 μm) were used in this study. Table 1 shows the target constitution of the four types of composites. In view of the brittle characteristics of Ni3Al at room temperature, when the Al fraction is more than 25 at.%, a non-stoichiometric nickel aluminide compound, Ni76Al24, can be widely used in practical applications [25]. Therefore, the target composition of the nickel aluminide was determined to be Ni76Al24 and 0.1 wt.% B was also added to decrease the
brittleness of the grain boundaries in the nickel aluminide. The weighed powders were placed into stainless steel vials with WC-Co balls and the vials were then filled with pure argon (99.999%). The mixtures were milled using planetary balls with a ball-to-powder weight ratio of 30:1 and a rotation speed of 400 rpm. The external temperature of the milling vial was monitored using an infrared thermometer (F62Max, Fluke, China) during milling to measure the ignition time of MSR.
Fig. 2. Morphology of the TiC–Ni3Al composite powder and EDS map scanning of metallic elements.
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After ignition, the powders continued to be milled for 1 h to complete the reaction and ensure the formation of homogenous products. The synthesized composite powders were sieved and compacted in a rectangular mold under a uniaxial pressure of 300 MPa. Then, the green compacts were sintered at 1415–1460 °C under vacuum (10−3– 10−2 Pa). The phases of the milled products were identified by X-ray diffraction (XRD-7000, Shimadzu, Japan). The morphology of the milled powders and microstructure of the sintered bulks were examined by Scanning Electron Microscope (SEM, Nova Nano 450, FEI, USA) equipped with an Energy Dispersive Spectrometer (EDS, Oxford, UK). The finer microstructure and interfacial structure of the sintered bulks were observed by STEM high-angle annular dark-field (HAADF) and HRTEM using a Tecnai G2F30 (FEI, USA). The porosity of the sintered bulks was evaluated using ISO 4505. The hardness was measured using a Vickers hardness tester (432 SVD, Wolpert Wilson Instrument, China) with an indenter load of 30 kg over 15 s, and fracture toughness (KIC) was evaluated by the cracks resulting from the indentation using the expression derived by Shetty [26]. The transverse rupture strength (TRS) was measured using the three-point bending method (the strain rate was 0.5 mm/min) by a universal material testing machine (Zwick/ Roell Z020, Germany), and the stress–strain curve was recorded. The dimensions of specimen were 20.00 mm × 6.50 mm × 5.00 mm (span 14.5 mm). 3. Results and discussion 3.1. Synthesis of TiC–Ni3Al composite powders An abrupt increase in the temperature of the milling vials was observed in the four Ti–C–Ni–Al mixtures after extensive milling time, which indicated that MSR had occurred in these systems. The ignition time, i.e. activation period, for each system is shown in Table 1 and this activation period lengthened with the increase of the binder
Fig. 4. XRD patterns of sintered TiC–xNi3Al composites.
content in the mixture. The phase composition of the powders following the MSR process is shown in Fig. 1(a). The resulting phases consisted solely of TiC and Ni3Al due to the high stability of these two materials among the possible compounds in the Ti–C–Ni–Al system (ΔHTiC = − 184.2 kJ/mol; ΔHNi3Al = − 153.4 J/mol; ΔHTiAl3 = −142.3 kJ/mol; ΔHNi3Ti = − 138.9 kJ/mol; ΔHTiAl = − 72.8 kJ/mol; ΔHNiAl = − 117.7 kJ/mol; ΔHNiTi = − 66.5 kJ/mol). Moreover, no Ni remnants or by-products such as NiAl were present, which usually appears in the synthesis of Ni3Al using SHS. This result indicated that the MSR process is more efficient than SHS in producing homogenous products. Table 1 shows the lattice constant of the Ni3Al products which calculated using the XRD data. These values were slightly higher than the
Fig. 3. SEM images of the blocks formed in MSR: (a) morphology of the blocks; (b) microstructure of the block in P2; (c) microstructure of the block in P3; (d) EDS analysis for the binder in P2 (only metallic element quantified data was given).
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Fig. 5. SEM-BSE images of sintered TiC–xNi3Al composites: (a) x = 20; (b) x = 30; (c) x = 40; (d) x = 50.
lattice constant of the Ni3Al provided by The Joint Committee on Powder Diffraction Standards — International Centre for Diffraction Data (JCPDS-ICDD), which was found to be 0.3560–0.3572 nm. In addition to the powders, the products in the four systems also contained some 1–20 mm irregular blocks, which had the same phase composition as the powders (Fig. 1(b)). Additionally, WC abrasives from the milling balls appeared in the powders milled for longer periods. The XRD pattern of sample P5 after 5 h of milling is shown in Fig. 1(a). Sample P5 consisted primarily of Ni and only a small amount of Ni3Al was formed. This implied that Ni3Al cannot be synthesized by
MSR via ball milling of Ni and Al, although it can be easily obtained by SHS [27–29]. This variation may be due to the greater heat dissipation of MSR that occurs in the SHS process. The small quantity of Ni3Al in P5 was produced by a gradual diffusion mechanism via high-energy ball milling. The similarity in lattice parameters between Ni3Al and Ni (JCPDS-ICDD provides the lattice parameter of Ni as 0.3524– 0.3540 nm) resulted in an overlap of their diffraction peaks in XRD pattern, especially when the particle size was reduced by milling. The above results suggested that the MSR that occurred in Ti–C–Ni–Al system can be attributed to the higher exothermic reaction between Ti
Fig. 6. STEM image of the TiC–Ni3Al composite: (a) core-rim structure; (b) EDS line analysis from A to B.
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and C (Tad(TiC) = 3210 K [30]) rather than the lower exothermic reactions between Ni and Al (Tad(Ni3Al) = 2079 K [31]). The morphology of the synthesized composite powder and the EDS scan of metallic elements are shown in Fig. 2. The composite powder had a homogenous chemical composition and consisted of ultrafine particles and agglomerates of nanoparticles. Fig. 3 shows the morphology and microstructure of the blocks formed in the MSR. Many cavities were observed in the blocks and a significant amount of spherical ceramic particles appeared inside these cavities (Fig. 3(a)). This morphology indicated that a liquid phase had formed during the MSR triggered by the ball milling [32]. The microstructure of the blocks formed in samples P2 and P3 (Fig. 3(b) and (c)) showed that TiC grains were distributed uniformly in the Ni3Al binder phase and the grain size in the blocks decreased with the increase of the Ni3Al content. The EDS analysis shown in Fig. 3(d) verified that the greater lattice parameters of the Ni3Al formed during the MSR resulted from the Ti that dissolved in the Ni3Al. It is worth noting that the microstructure of the blocks was very similar to the microstructure of the TiC–Ni and TiC–Ni3Al composites prepared using SHS [13,33]. In these reported researches, it has been verified that TiC was formed via dissolution–precipitation during the SHS process. In light of the dissolution of Ti in Ni3Al and the similarities between MSR and SHS, it is reasonable to conclude that the TiC particles during MSR also resulted from the dissolution–precipitation mechanism. 3.2. Microstructure of TiC–Ni3Al composites The XRD patterns and microstructure of the sintered TiC–Ni3Al composites with varying Ni3Al content are shown in Figs. 4 and 5. The XRD patterns suggest that the phase composition of the powders did not change during sintering. With the increase of Ni3Al, the volume of the liquid phase formed during sintering increased and additional liquid filled the space between the ceramic particles. Consequently, porosity of sintered bulks decreased and ceramic grains became refined, because the coalescence growth mechanism of the ceramic grains was inhibited by the molten binder. However, there was no drastic decrease in particle size when the Ni3Al binder was increased from 40 wt.% to 50 wt.% (Fig. 5(c) and (d)). It is well known that the ceramic grains grow by coalescence and dissolution–precipitation mechanisms during liquid phase sintering. When the volume of the liquid phase reaches a specific limit, coalescence growth is inhibited, which enhances the grain growth via the dissolution–precipitation mechanism. A core-rim structure grains are not easily observed in the SEM images, unless STEM-HAADF image procedure is conducted. The result is shown in Fig. 6(a), where the ceramic grains are composed of a gray core, a white inner rim near the core and a gray outer rim near the
Fig. 7. HRTEM image of the interface between the ceramic phase and the binder phase.
Table 2 Mechanical properties and porosity of TiC–Ni3Al composites prepared at the optimized sintering temperature. Composite
Hardness (GPa)
TRS (MPa)
KIC (MPa m1/2)
Porosity
TiC–20Ni3Al (1460 °C) TiC–30Ni3Al (1460 °C) TiC–40Ni3Al (1445 °C) TiC–50Ni3Al (1430 °C)
14.8 ± 0.3 12.7 ± 0.2 11.6 ± 0.2 9.8 ± 0.2
1220 ± 66 1428 ± 49 1735 ± 41 2165 ± 53
8.7 ± 0.9 10.7 ± 0.5 14.9 ± 0.3 17.5 ± 0.3
A06B02 A02B02 A02B02 A02B00
binder. The EDS line analysis from point A to B (Fig. 6(b)) indicated that the inner rim contained the heavy element W, which displayed a bright color in the STEM-HAADF image coinciding with the concentration of the heavy elements in the sample. From the XRD patterns and the EDS line analysis, it can be concluded that the sample cores were composed of undissolved TiC, the inner rims were (Ti,W)C solid solution and the outer rims were also composed of TiC. The (Ti,W)C solid solution cannot be identified by the XRD patterns as a result of a crystal structure and lattice parameters similar to TiC. The formation mechanism of this type of core-rim structure was related to the dissolution and precipitation behaviors of the carbides in the TiC–Ni3Al systems. During sintering, WC abrasives and partial TiC particles (smaller TiC particles dissolved preferentially due to their larger surface energy) dissolved continuously into the binder Ni3Al. These dissolved carbide constituents precipitated onto the surface of the undissolved TiC in the form of a (Ti,W)C solid solution when the binder was saturated and the rims were subsequently formed. In Ti(C,N)–WC–Ni cermet system, it is believed that the inner rims are formed by solid-state reactions during an early stage of sintering (~1300 °C), while the outer rims are formed during the subsequent liquid phase sintering and cooling stages [34,35]. The dissolution rate ratio of WC/Ti(C,N) in the binder at the early sintering stage is higher than at the late sintering stage [35]. This gives rise to the high W concentration in the inner rims. This probably also constituted the formation mechanism of the inner and outer rims in the TiC–Ni3Al systems. However, the very low WC content resulting from the milling balls led to the extremely thin inner rims, making them difficult to observe in SEM-BSE images. In conventional TiC and Ti(C,N)-based cermets incorporate secondary carbides such as Mo2C and WC, thereby, the outer rims contain substantial W and Mo due to the dissolution and precipitation of these carbides. This makes the surface energy between the rims and the binder become isotropic [35]. Consequently, the grains are round-edge
Fig. 8. Dislocation lines throughout the large ceramic grains of the TiC–Ni3Al composite.
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Fig. 9. Fracture morphology of the composites after TRS testing: (a) TiC–20Ni3Al; (b) TiC–30Ni3Al. The circled parts are tear ridges and the arrows indicate debonding of the ceramic grains.
shape in conventional TiC and Ti(C,N)-based cermets [35]. However, the grains were sharp-edged or faceted in the TiC–Ni3Al composites. Since W existed predominantly in the inner rims, only Ti precipitated in the form of TiC at the outer rim formation stage. Precipitated TiC grew continuously along with a stable crystal plane of TiC (for example (200) plane as shown in Fig. 7) by the Ostwald ripening mechanism during liquid phase sintering and gradually evolved into a faceted octahedron in the 3D space [36]. The HRTEM image of the interface (Fig. 7) showed a low-angle misorientation between the ceramic phase and the binder phase. Furthermore, the interface contained a 1–2 nm width of a transition zone where the ceramic phase and the binder phase shared atoms. 3.3. Mechanical properties of TiC–Ni3Al composites The optimum sintering temperature of the TiC–Ni3Al composites, their mechanical properties and the porosity obtained at the optimum sintering temperature are shown in Table 2. As can be seen, the TRS and KIC were dramatically enhanced by the increase in Ni3Al as a result
of the increase of the binder, whereas the composite hardness decreased due to the absence of a hard ceramic phase. Examination of the STEM image shown in Fig. 8, shows that dislocations appeared in the large ceramic grains, but were absent in small ceramic grains, which directly led to fracture of large grains when the material experienced an external stress. Consequently, a large number of splits and traces of cleavage in the ceramic phase can be observed in the fracture morphology of the TiC–Ni3Al composites containing low binder content (Fig. 9(a)). In these samples, typical transgranular fracture characters can be seen that were caused by coarse ceramic grains resulting in low TRS and KIC values. As the binder content in the sample was increased, the ceramic grains became refined. Evidence of splitting and cleavage traces in the ceramic phase of these samples were decreased, while traces of ceramic grains debonding (white arrows in Fig. 9(b)) and tear ridges induced by plastic deformation of the binder phase (circled in Fig. 9(b)) appeared. This failure manner indicated that the sample's fracture mode had changed to intergranular fracture, which resulted in improved TRS and KIC values.
Fig. 10. The stress–strain curves of the TiC–xNi3Al composites.
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From the properties listed in Table 2 and the stress–strain curves shown in Fig. 10, it can be seen that the TRS of the TiC–Ni3Al composites in this study was superior to that prepared using the MIS method [1] and by conventional vacuum sintering [7] (the TRS values resulting from these other approaches were less than 1200 MPa at room temperature). It is well known that the TRS is very sensitive to the presence of pores in a sintered body. Generally speaking, commercially-available TiC powder is used as the raw material in MIS and vacuum sintering methods. The surfaces of these raw powders absorb large quantities of oxygen and other contaminations. The oxygen content of the powder is further increased during the mixing process. This is harmful to the wettability of the liquid binder on the ceramic phase during sintering and inhibits the densification of the TiC–Ni3Al composites. Apparently, the improvement of the properties of the TiC–Ni3Al composites prepared in this study was a direct result of the special preparation methods used to treat the TiC–Ni3Al powders. This process may well have cleaned the interface between the TiC grains and the Ni3Al binder due to the dissolution and precipitation formation mechanisms of TiC particles during MSR. This may well have resulted in improved wettability between these two materials. 4. Conclusions Ultrafine TiC–Ni3Al composite powders with various binder contents were synthesized using the MSR method employing high energy ball milling of Ti–C–Ni–Al powder mixtures. MSR that took place in these systems was triggered by the high exothermic Ti–C reaction, and ball milling of Ni–Al mixture cannot generate self-sustained reaction. The heat released during MSR melted metallic powders and TiC particles were formed via dissolution and precipitation mechanisms. As Ni3Al content in the mixtures was increased, the porosity of the sintered TiC–Ni3Al composites decreased and the ceramic grains became refined. The TiC–Ni3Al composites possessed better greater mechanical properties than similar materials prepared using MIS and conventional vacuum sintering methods, indicating that MSR is a promising method for preparing practical TiC–Ni3Al materials. Acknowledgments This research is funded by the National Natural Science Foundation of China (51375195) and Science and Technology Plan Projects of Shenzhen City (JCYJ20140903171444755). The authors thank the Analytical and Testing Center of Huazhong University of Science and Technology. References [1] P.F. Becher, K.P. Plucknett, Properties of Ni3Al-bonded titanium carbide ceramics, J. Eur. Ceram. Soc. 18 (1998) 395–400. [2] L. Chen, W. Wen, H. Cui, Yielding description for a Ni3Al based intermetallic alloy, Mater. Des. 41 (2012) 192–197. [3] N.S. Stoloff, C.T. Liu, S.C. Deevi, Emerging applications of intermetallics, Intermetallics 8 (2000) 1313–1320. [4] J.W. Cohron, E.P. George, L. Heatherly, C.T. Liu, R.H. Zee, Reply to "a comment on hydrogen–boron interaction and its effect on the ductility and fracture of Ni3Al", Scripta Mater. 38 (1998) 847–850. [5] J.W. Cohron, E.P. George, L. Heatherly, C.T. Liu, R.H. Zee, Hydrogen–boron interaction and its effect on the ductility and fracture of Ni3Al, Acta Mater. 45 (1997) 2801–2811. [6] E.M. Schulson, Y. Xu, On the ductility of Ni3Al: effects of strain rate, environment and boron, Acta Mater. 45 (1997) 3491–3494.
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Materials and Design journal homepage: www.elsevier.com/locate/jmad
Characterization of intermetallic bonded TiC composites prepared by mechanically induced self-sustained reaction Xiao Chen a, Jianfeng Xu a,⁎, Quanquan Sun a, Wei Yang a, Weihao Xiong b a b
State Key Laboratory of Digital Manufacturing Equipment & Technology, School of Mechanical Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China
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Article history: Received 30 June 2015 Received in revised form 27 September 2015 Accepted 28 September 2015 Available online xxxx Keywords: TiC–Ni3Al Mechanically induced self-sustained reaction Microstructure Mechanical properties
a b s t r a c t The microstructure and mechanical properties of Ni3Al bonded TiC composites prepared by mechanically induced self-sustained reaction (MSR) were investigated. Ultrafine TiC–Ni3Al composite powders with various binder contents were synthesized by high energy ball milling of Ti–C–Ni–Al powder mixtures. Some porous blocks, which consisted of spherical TiC grains and Ni3Al binder, were also formed in each system after MSR. A core-rim structure present in the grains of the sintered TiC–Ni3Al composites was observed by Scanning Transmission Electron Microscopy (STEM). TiC particles dissolved selectively and then precipitated along with a stable crystal plane and gradually evolved into a faceted octahedron in the 3D space, which led to the formation of sharp-edged ceramic grains. High Resolution Transmission Electron Microscopy (HRTEM) analysis showed a 1–2 nm width transition zone and a low-angle misorientation between the ceramic phase and the binder phase. The TiC particles that precipitated from Ni3Al during MSR were easily wetted by the liquid Ni3Al binder, which gave rise to the superior mechanical properties of the TiC–Ni3Al composites. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Nickel aluminide (Ni3Al) has attracted considerable attention as an alternative binder for ceramic matrix composites due to its exceptional high temperature strength and excellent chemical stability [1–3]. The brittleness of Ni3Al appears to have been solved by doping boron [4– 6] and the good wettability of Ni3Al on TiC [7,8] makes it possible to use Ni3Al as a binder for TiC-based composites. To date, extensive attempts have been made to prepare TiC–Ni3Al composites [1,9–16]. This type of Ni3Al bonded TiC composites exhibit good oxidation and corrosion resistance, as well as high strength at elevated temperatures [1,16]. As such, they are promising materials for applications in high temperature, corrosion and wear environments, including cutting tools, metal heat forming tools and turbine blades for gas-turbine engines. However, the inferior flexure strength of TiC–Ni3Al composites at room temperature limits their practical application. Melt-infiltration sintering (MIS) has been adopted to prepare TiC– Ni3Al composites [1,9,11,15–17]. Some researchers have also prepared TiC–Ni3Al composites using vacuum sintering [7,10], reactive sintering [14,18] and self-propagating high temperature synthesis (SHS) [12, 13]. Mechanically induced self-sustaining reactions (MSR) have been used as a powerful manufacturing technique for the synthesis of a great number of novel materials [19–24]. This reactive milling method takes advantage of the strong exothermic character of the formation ⁎ Corresponding author. E-mail address: [email protected] (J. Xu).
http://dx.doi.org/10.1016/j.matdes.2015.09.158 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
of compounds from chemical elements to promote self-propagating reactions during milling and obtain particulate materials with high purity and homogenous composition. MSR has been shown to offer many advantages over other methods in terms of low cost, small particle sizes, and narrow size distribution products [20]. Moreover, the structure of these particles steadily diminishes and many dislocations are created during high-energy milling, which improves the mechanical properties of these novel materials [20,21]. In this reported work, various TiC– Ni3Al composites with binder contents from 20 wt.% to 50 wt.% were prepared using MSR with subsequent vacuum sintering. The synthesis of composite powders together with the microstructure and mechanical properties of the TiC–xNi3Al composites was investigated. 2. Experimental procedures Ti powder (99.0% pure, 45 μm), Ni powder (99.8% pure, 2.25 μm), Al powder (99.8% pure, 1.80 μm), graphite powder (99.9% pure, 5.50 μm)
Table 1 Target compositions, ignition time and lattice parameter of the synthesized Ni3Al. Mixture
Target composition (wt.%)
Ignition time
Lattice parameter (nm)
P1 P2 P3 P4 P5
TiC–20Ni3Al TiC–30Ni3Al TiC–40Ni3Al TiC–50Ni3Al Ni3Al
93 min 121 min 155 min 184 min –
0.3608 0.3602 0.3599 0.3594 0.3582
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Fig. 1. XRD patterns of the powders and blocks formed in Ti–C–Ni–Al system after MSR: (a) the powders formed in P1–P5; (b) the blocks formed in P1–P4.
and B powder (99.0% pure, 1–15 μm) were used in this study. Table 1 shows the target constitution of the four types of composites. In view of the brittle characteristics of Ni3Al at room temperature, when the Al fraction is more than 25 at.%, a non-stoichiometric nickel aluminide compound, Ni76Al24, can be widely used in practical applications [25]. Therefore, the target composition of the nickel aluminide was determined to be Ni76Al24 and 0.1 wt.% B was also added to decrease the
brittleness of the grain boundaries in the nickel aluminide. The weighed powders were placed into stainless steel vials with WC-Co balls and the vials were then filled with pure argon (99.999%). The mixtures were milled using planetary balls with a ball-to-powder weight ratio of 30:1 and a rotation speed of 400 rpm. The external temperature of the milling vial was monitored using an infrared thermometer (F62Max, Fluke, China) during milling to measure the ignition time of MSR.
Fig. 2. Morphology of the TiC–Ni3Al composite powder and EDS map scanning of metallic elements.
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After ignition, the powders continued to be milled for 1 h to complete the reaction and ensure the formation of homogenous products. The synthesized composite powders were sieved and compacted in a rectangular mold under a uniaxial pressure of 300 MPa. Then, the green compacts were sintered at 1415–1460 °C under vacuum (10−3– 10−2 Pa). The phases of the milled products were identified by X-ray diffraction (XRD-7000, Shimadzu, Japan). The morphology of the milled powders and microstructure of the sintered bulks were examined by Scanning Electron Microscope (SEM, Nova Nano 450, FEI, USA) equipped with an Energy Dispersive Spectrometer (EDS, Oxford, UK). The finer microstructure and interfacial structure of the sintered bulks were observed by STEM high-angle annular dark-field (HAADF) and HRTEM using a Tecnai G2F30 (FEI, USA). The porosity of the sintered bulks was evaluated using ISO 4505. The hardness was measured using a Vickers hardness tester (432 SVD, Wolpert Wilson Instrument, China) with an indenter load of 30 kg over 15 s, and fracture toughness (KIC) was evaluated by the cracks resulting from the indentation using the expression derived by Shetty [26]. The transverse rupture strength (TRS) was measured using the three-point bending method (the strain rate was 0.5 mm/min) by a universal material testing machine (Zwick/ Roell Z020, Germany), and the stress–strain curve was recorded. The dimensions of specimen were 20.00 mm × 6.50 mm × 5.00 mm (span 14.5 mm). 3. Results and discussion 3.1. Synthesis of TiC–Ni3Al composite powders An abrupt increase in the temperature of the milling vials was observed in the four Ti–C–Ni–Al mixtures after extensive milling time, which indicated that MSR had occurred in these systems. The ignition time, i.e. activation period, for each system is shown in Table 1 and this activation period lengthened with the increase of the binder
Fig. 4. XRD patterns of sintered TiC–xNi3Al composites.
content in the mixture. The phase composition of the powders following the MSR process is shown in Fig. 1(a). The resulting phases consisted solely of TiC and Ni3Al due to the high stability of these two materials among the possible compounds in the Ti–C–Ni–Al system (ΔHTiC = − 184.2 kJ/mol; ΔHNi3Al = − 153.4 J/mol; ΔHTiAl3 = −142.3 kJ/mol; ΔHNi3Ti = − 138.9 kJ/mol; ΔHTiAl = − 72.8 kJ/mol; ΔHNiAl = − 117.7 kJ/mol; ΔHNiTi = − 66.5 kJ/mol). Moreover, no Ni remnants or by-products such as NiAl were present, which usually appears in the synthesis of Ni3Al using SHS. This result indicated that the MSR process is more efficient than SHS in producing homogenous products. Table 1 shows the lattice constant of the Ni3Al products which calculated using the XRD data. These values were slightly higher than the
Fig. 3. SEM images of the blocks formed in MSR: (a) morphology of the blocks; (b) microstructure of the block in P2; (c) microstructure of the block in P3; (d) EDS analysis for the binder in P2 (only metallic element quantified data was given).
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Fig. 5. SEM-BSE images of sintered TiC–xNi3Al composites: (a) x = 20; (b) x = 30; (c) x = 40; (d) x = 50.
lattice constant of the Ni3Al provided by The Joint Committee on Powder Diffraction Standards — International Centre for Diffraction Data (JCPDS-ICDD), which was found to be 0.3560–0.3572 nm. In addition to the powders, the products in the four systems also contained some 1–20 mm irregular blocks, which had the same phase composition as the powders (Fig. 1(b)). Additionally, WC abrasives from the milling balls appeared in the powders milled for longer periods. The XRD pattern of sample P5 after 5 h of milling is shown in Fig. 1(a). Sample P5 consisted primarily of Ni and only a small amount of Ni3Al was formed. This implied that Ni3Al cannot be synthesized by
MSR via ball milling of Ni and Al, although it can be easily obtained by SHS [27–29]. This variation may be due to the greater heat dissipation of MSR that occurs in the SHS process. The small quantity of Ni3Al in P5 was produced by a gradual diffusion mechanism via high-energy ball milling. The similarity in lattice parameters between Ni3Al and Ni (JCPDS-ICDD provides the lattice parameter of Ni as 0.3524– 0.3540 nm) resulted in an overlap of their diffraction peaks in XRD pattern, especially when the particle size was reduced by milling. The above results suggested that the MSR that occurred in Ti–C–Ni–Al system can be attributed to the higher exothermic reaction between Ti
Fig. 6. STEM image of the TiC–Ni3Al composite: (a) core-rim structure; (b) EDS line analysis from A to B.
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and C (Tad(TiC) = 3210 K [30]) rather than the lower exothermic reactions between Ni and Al (Tad(Ni3Al) = 2079 K [31]). The morphology of the synthesized composite powder and the EDS scan of metallic elements are shown in Fig. 2. The composite powder had a homogenous chemical composition and consisted of ultrafine particles and agglomerates of nanoparticles. Fig. 3 shows the morphology and microstructure of the blocks formed in the MSR. Many cavities were observed in the blocks and a significant amount of spherical ceramic particles appeared inside these cavities (Fig. 3(a)). This morphology indicated that a liquid phase had formed during the MSR triggered by the ball milling [32]. The microstructure of the blocks formed in samples P2 and P3 (Fig. 3(b) and (c)) showed that TiC grains were distributed uniformly in the Ni3Al binder phase and the grain size in the blocks decreased with the increase of the Ni3Al content. The EDS analysis shown in Fig. 3(d) verified that the greater lattice parameters of the Ni3Al formed during the MSR resulted from the Ti that dissolved in the Ni3Al. It is worth noting that the microstructure of the blocks was very similar to the microstructure of the TiC–Ni and TiC–Ni3Al composites prepared using SHS [13,33]. In these reported researches, it has been verified that TiC was formed via dissolution–precipitation during the SHS process. In light of the dissolution of Ti in Ni3Al and the similarities between MSR and SHS, it is reasonable to conclude that the TiC particles during MSR also resulted from the dissolution–precipitation mechanism. 3.2. Microstructure of TiC–Ni3Al composites The XRD patterns and microstructure of the sintered TiC–Ni3Al composites with varying Ni3Al content are shown in Figs. 4 and 5. The XRD patterns suggest that the phase composition of the powders did not change during sintering. With the increase of Ni3Al, the volume of the liquid phase formed during sintering increased and additional liquid filled the space between the ceramic particles. Consequently, porosity of sintered bulks decreased and ceramic grains became refined, because the coalescence growth mechanism of the ceramic grains was inhibited by the molten binder. However, there was no drastic decrease in particle size when the Ni3Al binder was increased from 40 wt.% to 50 wt.% (Fig. 5(c) and (d)). It is well known that the ceramic grains grow by coalescence and dissolution–precipitation mechanisms during liquid phase sintering. When the volume of the liquid phase reaches a specific limit, coalescence growth is inhibited, which enhances the grain growth via the dissolution–precipitation mechanism. A core-rim structure grains are not easily observed in the SEM images, unless STEM-HAADF image procedure is conducted. The result is shown in Fig. 6(a), where the ceramic grains are composed of a gray core, a white inner rim near the core and a gray outer rim near the
Fig. 7. HRTEM image of the interface between the ceramic phase and the binder phase.
Table 2 Mechanical properties and porosity of TiC–Ni3Al composites prepared at the optimized sintering temperature. Composite
Hardness (GPa)
TRS (MPa)
KIC (MPa m1/2)
Porosity
TiC–20Ni3Al (1460 °C) TiC–30Ni3Al (1460 °C) TiC–40Ni3Al (1445 °C) TiC–50Ni3Al (1430 °C)
14.8 ± 0.3 12.7 ± 0.2 11.6 ± 0.2 9.8 ± 0.2
1220 ± 66 1428 ± 49 1735 ± 41 2165 ± 53
8.7 ± 0.9 10.7 ± 0.5 14.9 ± 0.3 17.5 ± 0.3
A06B02 A02B02 A02B02 A02B00
binder. The EDS line analysis from point A to B (Fig. 6(b)) indicated that the inner rim contained the heavy element W, which displayed a bright color in the STEM-HAADF image coinciding with the concentration of the heavy elements in the sample. From the XRD patterns and the EDS line analysis, it can be concluded that the sample cores were composed of undissolved TiC, the inner rims were (Ti,W)C solid solution and the outer rims were also composed of TiC. The (Ti,W)C solid solution cannot be identified by the XRD patterns as a result of a crystal structure and lattice parameters similar to TiC. The formation mechanism of this type of core-rim structure was related to the dissolution and precipitation behaviors of the carbides in the TiC–Ni3Al systems. During sintering, WC abrasives and partial TiC particles (smaller TiC particles dissolved preferentially due to their larger surface energy) dissolved continuously into the binder Ni3Al. These dissolved carbide constituents precipitated onto the surface of the undissolved TiC in the form of a (Ti,W)C solid solution when the binder was saturated and the rims were subsequently formed. In Ti(C,N)–WC–Ni cermet system, it is believed that the inner rims are formed by solid-state reactions during an early stage of sintering (~1300 °C), while the outer rims are formed during the subsequent liquid phase sintering and cooling stages [34,35]. The dissolution rate ratio of WC/Ti(C,N) in the binder at the early sintering stage is higher than at the late sintering stage [35]. This gives rise to the high W concentration in the inner rims. This probably also constituted the formation mechanism of the inner and outer rims in the TiC–Ni3Al systems. However, the very low WC content resulting from the milling balls led to the extremely thin inner rims, making them difficult to observe in SEM-BSE images. In conventional TiC and Ti(C,N)-based cermets incorporate secondary carbides such as Mo2C and WC, thereby, the outer rims contain substantial W and Mo due to the dissolution and precipitation of these carbides. This makes the surface energy between the rims and the binder become isotropic [35]. Consequently, the grains are round-edge
Fig. 8. Dislocation lines throughout the large ceramic grains of the TiC–Ni3Al composite.
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Fig. 9. Fracture morphology of the composites after TRS testing: (a) TiC–20Ni3Al; (b) TiC–30Ni3Al. The circled parts are tear ridges and the arrows indicate debonding of the ceramic grains.
shape in conventional TiC and Ti(C,N)-based cermets [35]. However, the grains were sharp-edged or faceted in the TiC–Ni3Al composites. Since W existed predominantly in the inner rims, only Ti precipitated in the form of TiC at the outer rim formation stage. Precipitated TiC grew continuously along with a stable crystal plane of TiC (for example (200) plane as shown in Fig. 7) by the Ostwald ripening mechanism during liquid phase sintering and gradually evolved into a faceted octahedron in the 3D space [36]. The HRTEM image of the interface (Fig. 7) showed a low-angle misorientation between the ceramic phase and the binder phase. Furthermore, the interface contained a 1–2 nm width of a transition zone where the ceramic phase and the binder phase shared atoms. 3.3. Mechanical properties of TiC–Ni3Al composites The optimum sintering temperature of the TiC–Ni3Al composites, their mechanical properties and the porosity obtained at the optimum sintering temperature are shown in Table 2. As can be seen, the TRS and KIC were dramatically enhanced by the increase in Ni3Al as a result
of the increase of the binder, whereas the composite hardness decreased due to the absence of a hard ceramic phase. Examination of the STEM image shown in Fig. 8, shows that dislocations appeared in the large ceramic grains, but were absent in small ceramic grains, which directly led to fracture of large grains when the material experienced an external stress. Consequently, a large number of splits and traces of cleavage in the ceramic phase can be observed in the fracture morphology of the TiC–Ni3Al composites containing low binder content (Fig. 9(a)). In these samples, typical transgranular fracture characters can be seen that were caused by coarse ceramic grains resulting in low TRS and KIC values. As the binder content in the sample was increased, the ceramic grains became refined. Evidence of splitting and cleavage traces in the ceramic phase of these samples were decreased, while traces of ceramic grains debonding (white arrows in Fig. 9(b)) and tear ridges induced by plastic deformation of the binder phase (circled in Fig. 9(b)) appeared. This failure manner indicated that the sample's fracture mode had changed to intergranular fracture, which resulted in improved TRS and KIC values.
Fig. 10. The stress–strain curves of the TiC–xNi3Al composites.
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From the properties listed in Table 2 and the stress–strain curves shown in Fig. 10, it can be seen that the TRS of the TiC–Ni3Al composites in this study was superior to that prepared using the MIS method [1] and by conventional vacuum sintering [7] (the TRS values resulting from these other approaches were less than 1200 MPa at room temperature). It is well known that the TRS is very sensitive to the presence of pores in a sintered body. Generally speaking, commercially-available TiC powder is used as the raw material in MIS and vacuum sintering methods. The surfaces of these raw powders absorb large quantities of oxygen and other contaminations. The oxygen content of the powder is further increased during the mixing process. This is harmful to the wettability of the liquid binder on the ceramic phase during sintering and inhibits the densification of the TiC–Ni3Al composites. Apparently, the improvement of the properties of the TiC–Ni3Al composites prepared in this study was a direct result of the special preparation methods used to treat the TiC–Ni3Al powders. This process may well have cleaned the interface between the TiC grains and the Ni3Al binder due to the dissolution and precipitation formation mechanisms of TiC particles during MSR. This may well have resulted in improved wettability between these two materials. 4. Conclusions Ultrafine TiC–Ni3Al composite powders with various binder contents were synthesized using the MSR method employing high energy ball milling of Ti–C–Ni–Al powder mixtures. MSR that took place in these systems was triggered by the high exothermic Ti–C reaction, and ball milling of Ni–Al mixture cannot generate self-sustained reaction. The heat released during MSR melted metallic powders and TiC particles were formed via dissolution and precipitation mechanisms. As Ni3Al content in the mixtures was increased, the porosity of the sintered TiC–Ni3Al composites decreased and the ceramic grains became refined. The TiC–Ni3Al composites possessed better greater mechanical properties than similar materials prepared using MIS and conventional vacuum sintering methods, indicating that MSR is a promising method for preparing practical TiC–Ni3Al materials. Acknowledgments This research is funded by the National Natural Science Foundation of China (51375195) and Science and Technology Plan Projects of Shenzhen City (JCYJ20140903171444755). The authors thank the Analytical and Testing Center of Huazhong University of Science and Technology. References [1] P.F. Becher, K.P. Plucknett, Properties of Ni3Al-bonded titanium carbide ceramics, J. Eur. Ceram. Soc. 18 (1998) 395–400. [2] L. Chen, W. Wen, H. Cui, Yielding description for a Ni3Al based intermetallic alloy, Mater. Des. 41 (2012) 192–197. [3] N.S. Stoloff, C.T. Liu, S.C. Deevi, Emerging applications of intermetallics, Intermetallics 8 (2000) 1313–1320. [4] J.W. Cohron, E.P. George, L. Heatherly, C.T. Liu, R.H. Zee, Reply to "a comment on hydrogen–boron interaction and its effect on the ductility and fracture of Ni3Al", Scripta Mater. 38 (1998) 847–850. [5] J.W. Cohron, E.P. George, L. Heatherly, C.T. Liu, R.H. Zee, Hydrogen–boron interaction and its effect on the ductility and fracture of Ni3Al, Acta Mater. 45 (1997) 2801–2811. [6] E.M. Schulson, Y. Xu, On the ductility of Ni3Al: effects of strain rate, environment and boron, Acta Mater. 45 (1997) 3491–3494.
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