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Al2O3 composites
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Effects of TiN addition on the properties of hot pressed TiN–Ti/Al2O3 composites Qian Liu, Fangfang Qi, Junyan Wu, Liu Zhang, Guopu Shi, Zhi Wang
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School of Material Science and Engineering, University of Jinan, Jinan 250022, China
A R T I C LE I N FO
A B S T R A C T
Keywords: TiN Resistivity Cermet composite Vacuum hot-pressing sintering
TiN–Ti/Al2O3 composites of varying TiN content (0–20 vol%) were prepared by vacuum hot-pressing sintering at different temperatures (1400 °C and 1500 °C) to investigate how TiN affected the mechanical properties and electrical conductivity of the composites. Sintered samples with added TiN exhibited better performance than those without it. The sample with 20 vol% TiN sintered 1500 °C had an optimal relative density of 99.49, Vickers hardness of 14.94 GPa, flexural strength of 321.55 MPa, and electrical resistivity of 1474.7 μΩ cm. However, this increased temperature did not improve the best sample resistivity of 930.3 μΩ cm, which was obtained at 1400 °C. Form SEM images and XRD patterns, the positive effect of TiN on composite mechanical properties may be ascribed to its good performance of high hardness and strength, a decrease of the brittle intermetallic phase, the form of AlTi3N, and the impact of the fine-grained strength of the TiN phase.
1. Introduction Cermet materials have attracted wide attention because their properties originate from both metal and ceramic materials, and include wear and corrosion resistance, high toughness, and strength [1–4]. As one of the most commonly-used high-temperature structural ceramics, Al2O3 is widely used in the semi-conductor, aerospace, and military industries due to its high hardness and good resistance to high temperature, corrosion, and wear [5–7]. On the other hand, metal Ti has been used in the aerospace, marine, and medical fields since it is tough, ductile, and electrically-conductive [8,9]. Most importantly, the two materials are physically and chemically compatible and share similar thermal expansion coefficients. Therefore, they meet the requirements for preparing Ti/Al2O3 cermet, popular in aerodynamicengine and other automotive applications [10–12]. Unfortunately, in long-term studies of Ti/Al2O3 composites, intermetallic compounds of Ti–Al such as TiAl and Ti3Al are always found. These are characterized by low fracture toughness at room temperature and low oxidation resistance at elevated temperatures, significantly restricting the application of Ti/Al2O3 [13–17]. To solve this problem, interface reaction control in Ti–Al materials has been investigated in depth. D. Pilone [13] and Liu et al. [18] reported that Cr, Si, Nb, Y, and W can be alloyed with Ti–Al to reduce the production of Ti–Al compounds by generating other substances. Ai et al. [19] showed that Ti–Al composites could be reinforced by Al2O3 and Nb2O5. Compound Al-TiAl3 materials have also been prepared by
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uniform droplet spraying as described by Wang [14]. Excessive production intermetallic Ti–Al compounds can also be suppressed by adding metals such as Nb or metal oxides such as Y2O3 [20]. For instance, Wu et al. prepared laminar Ti/Al2O3 composites by casting method [21]. In addition, Nb, CeO2, and Y2O3 were doped to reduce the diffusion of atoms at the interface of Ti/Al2O3 composites to restrain the formation of intermetallic Ti–Al compounds on the basis of the lamellar material, by Liu et al. [22,23]. Clearly, despite much effort, the brittleness of intermetallic Ti–Al compounds remains an open issue. Under these circumstances, we find that TiN has long been considered to be ideal in the applications such cutting tools, wear-resistant materials, decorative materials and electrodes owing to its high hardness, beautiful golden color, high chemical stability, and good electrical conductivity [24–26]. Simultaneously, TiN is often used to toughen and strengthen Al2O3 because of its effective pinning. Thus, we hypothesize that when TiN is introduced into Ti/Al2O3 composites, TiN may not only toughen and enhance the composites, but also reduce the generation of Ti–Al intermetallic compounds by combination with N. To date, many studies have centered on TiN, but none has explicitly focused on its effect on TiN–Ti/Al2O3 composites. Therefore, the impact of TiN on the mechanical properties and microstructure of TiN–Ti/ Al2O3 is studied here, and the changes in conductivity and phase with increasing TiN are also characterized.
Corresponding author. E-mail address: [email protected] (Z. Wang).
https://doi.org/10.1016/j.ceramint.2018.03.125 Received 27 February 2018; Received in revised form 13 March 2018; Accepted 14 March 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Liu, Q., Ceramics International (2018), https://doi.org/10.1016/j.ceramint.2018.03.125
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Table 1 Chemical composition of group A samples (wt%). Samples
Ti (vol%)
Al2O3 (vol%)
TiN (vol%)
A0 A1 A2 A3 A4 A5 A6
40 40 40 40 40 40 –
60 58 56 54 52 50 60
– 2 4 6 8 10 40
Table 2 Chemical composition of group B samples (wt%). Samples
Ti (vol%)
Al2O3 (vol%)
TiN (vol%)
B0 B1 B2 B3 B4 B5
40 35 30 25 20 –
60 60 60 60 60 60
– 5 10 15 20 40
Fig. 1. XRD patterns of group A experiment.
2. Experimental procedures In this work, commercially available Al2O3 (99.9% pure, 1.5 µm), Ti (99.9% pure, 20 µm) and TiN (99.5% pure, 1 µm) from ST-NANO C o., Ltd. (Shanghai, China) were used as raw materials. In raw Al2O3 powder, α-Al2O3 accounts for the largest proportion. The experiment was divided into two groups, and the specimen design is summarized in Tables 1 and 2, respectively. Each group includes a TiN/Al2O3 sample in order to explore the reaction between Ti, TiN, and Al2O3. The TiN–Ti/Al2O3 composites were prepared as follows. After weighing, powders were mixed with a planetary ball mill (XQM-2, China) in an ethanol bath (99.7% pure) for 2 h at a milling speed of 200 rpm with a ball-to-powder weight ratio of 2:1. Then, the slurry was dried at 70 °C for 24 h in a vacuum oven. The dry powder was then placed into a graphite mold and sintered by vacuum hot-pressing (VVPgr-80-2300, Shanghai Haoyue, China). Sintering was performed with temperature rise of 10 °C/min (0–1200 °C), and 5 °C/min (1200–1500 °C), at which time the sintering temperature was held at 1400 °C (group A) and 1500 °C (group B) for 1.5 h under a load of 30 MPa. Lastly, sintered samples were removed when the temperature had naturally dropped to about 100 °C, after which the samples were cut, ground, and polished. The relative densities of the samples were determined by the Archimedes method. Using the three-point bending method, the flexural strength of the treated specimens was measured by an electromechanical testing machine (CMT5504, MTS Systems, China) at room temperature with a crosshead speed of 0.5 mm/min and span of 30 mm, with specimens sizes limited to 3 mm × 4 mm × 35 mm. The microhardness was obtained with a Vickers hardness tester (HV-10001S, China) at room temperature with a load of 1000 gf and indentation time was maintained for 15 s. Phase analysis of the samples was performed using an X-ray diffraction analyzer (D8-ADVANCE, Germany) with CuKα radiation, while the microstructure and elemental distributions of the composites were further analyzed with a scanning electron microscope (FEI QUANTA FEG 250, United States), and an electron probe micro-analyzer (Shimadzu EMPA-1600, Japan). In addition, the electrical resistivity was measured using the DC four-probe method with a current-reversal technique [27–29]. All characterization and tests were performed at room temperature.
Fig. 2. XRD patterns of group B experiment.
3. Results and discussion 3.1. Effect of TiN on phase composition of TiN–Ti/Al2O3 composites The results of X-ray diffraction analyses performed to determine the phase constitution of the TiN–Ti/Al2O3 composites are shown in Figs. 1 and 2. The new phase AlTi3N appeared with the addition of TiN, and the diffraction peaks of the TiN phases became stronger with increasing TiN content. No reactions occurred between TiN and Al2O3 for no other phase existed after sintering, as shown at points A6 and B5. As shown in Fig. 1 (A0–A5), the amount of Ti3O gradually reduced with increasing TiN content, because the oxygen diffusion was hindered by the addition of TiN. From A0 to A4, however, the formation of AlTi3N increased gradually, stabilizing at A4 and A5. Meanwhile, the amount of Ti3Al decreased when 2 vol% of TiN was added, as shown in Fig. 1 A1, remaining stable as additional TiN was added. From this, the production of AlTi3N requires two procedures: first, react Ti with Al2O3 to produce Ti3Al, then combine Ti3Al with N to generate AlTi3N. This process also is responsible for reducing the amount of Ti3Al. As shown in Fig. 2, in contrast to samples from group A samples, TiO was found in group B samples sintered at 1500 °C. The amount of TiO decreased until eventually becoming zero as the amount of TiN increased, in line with the trend of Ti3O. The reduction of TiO and Ti3O should be attributed to the decrease of Ti, and the hindering effect of TiN on the diffusion of O elements. In group B, Ti was reduced due to the increased TiN, lowering the Ti3Al. Moreover, a portion of the Ti3Al combined with N to generate AlTi3N. Therefore, the increased TiN led 2
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Fig. 3. SEM (SE) micrographs of the fracture surfaces of group A: A0, (b) A3, (c) A6.
Unfortunately, both caused weak mechanical performance of the composite. As is illustrated in Fig. 3b, when TiN was added to the composite, the phase distribution remained inhomogeneous, while the voids were drastically reduced as compared with the blank sample, and a good combination of phases appeared. From Fig. 3c, the distribution of phases was fairly uniform in the fractured TiN/Al2O3 composite. However, the sample was not visibly densified, because of the requirement for a higher sintering temperature for the TiN/Al2O3 composite. The relative densities of group A samples are listed in Table 3. Overall, the effect of the TiN phase on the density was negligible and the relative density rose slightly because of the reaction between Ti, TiN, and Al2O3. The fracture surfaces of group B samples were observed using SEM in the BSE mode, as shown in Fig. 4. The inhomogeneous phases were still untreated when the sintering temperature reached 1500 °C. However, this gradually improved with the addition of TiN, as shown in Fig. 4b–d, as the increased TiN content led to phase homogenize, especially for the TiN/Al2O3 composite. In addition, the Ti and Al2O3 grain sizes were refined by increasing the TiN content, due to its ability
Table 3 The relative densities of group A samples. Sample
A0
A1
A2
A3
A4
A5
A6
Density (%)
98.32
97.98
98.26
98.96
99.27
99.55
91.59
to a decrease in Ti3Al, but the amount of Ti3Al was too low to be detected when the content of TiN reached 15 vol%. At the same time, the amount of AlTi3N decreased owing to the decline of Ti3Al.
3.2. Effect of TiN on micro-structure of TiN–Ti/Al2O3 composites As shown in Fig. 3, the microstructure of TiN–Ti/Al2O3 cermets with different TiN content sintered at 1400 °C was observed by SEM in the SE mode. Bulk agglomerations are seen in Fig. 3a and massive voids are observed in the enlarged diagram on the right. This implies that Ti and Al2O3 were dispersed very unevenly in the blank sample lacking the TiN phase, while the blank sample A0 had a poor sintering density. 3
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Fig. 4. SEM (BSE) micrographs of the fracture surfaces of group B: (a) B0, (b) B2, (c) B4, (d) B5.
actual test, Ti and N corresponded well, and no separate Ti is shown in Fig. 6. The local Ti and N areas are enlarged and shown in Fig. 6b, indicating that more Ti was present than N despite the overlap between them. Because the amount of Ti was reduced and that of N increased, the local areas marked in the figure are considered to contribute to the TiN phase. Thus, Ti and TiN were mutually attractive during sintering and were mixed together in the sintered sample, reducing the formation of the intermetallic TiAl to some extent.
to pin particles. The bending strength can be enhanced by grain refinement according to the Hall-Petch formula,
σy = σ0 + Kd−1/2 where σy is the yield strength and d is the grain size. In addition, all samples sintered at 1500 °C had high densities with few voids, as shown in Table 4 where the relative densities of all samples were greater than 99%. Thus, the increased sintering temperature promoted the densification of the composites, while the effect of TiN was slight, in good agreement with the change seen in Fig. 4. Comparing Figs. 3c and 4d, the grains in TiN/Al2O3 sintered at 1500 °C were combined better. This confirms the conjecture that the TiN/Al2O3 composite required a higher sintering temperature than Ti/Al2O3. In Fig. 5a, the existence of a crystalline morphology and cracks along grain boundaries are shown, both suggesting that the fracture mode of the composites without added TiN was typically inter-granular. In contrast with B0, the cracks passing through grains in Fig. 5b demonstrate that the fracture mode of B4 with the added 20% TiN had changed from a single inter granular fracture to a mixture of inter granular and transgranular fracture. The change of the fracture mode is indicative of improved composite mechanical properties. In order to observe the distribution of the three main phases, the electron probe micro-analyzer was used to test a section of sample B4, with the results presented in Fig. 6. In Fig. 6a, bright areas are Ti and N element and dark areas are Al and O. In the latter, Al diffused differently from O promoting the formation of AlTi3N. From the XRD pattern, the main phase of the material comprised TiN, Ti, and Al2O3. In the
3.3. Effect of TiN on mechanical properties of TiN–Ti/Al2O3 composites In order to further explore the effects of TiN on Ti/Al2O3 composites, the bending strength and micro-hardness of samples in group A and B were measured, as shown in Figs. 7 and 8 respectively. From Fig. 7, the flexural strength and micro-hardness of group A were consistent. When the TiN content was less than 6 vol%, the decline of density was responsible for the decrease in bending strength and microhardness of samples in the earlier stage. The density increased with added TiN, accompanied by reinforced flexural strength and microhardness. Later, the improved density and weaker mechanical properties appeared simultaneously. This is because early in the addition procedure, the strength and hardness of group A samples depended mainly on their density, but afterwards, the amount of AlTi3N capable of promoting the combination of the three phases (Ti, TiN, Al2O3) did not continue to increase, causing the mechanical performance to degrade. At the same time, Al2O3 was substituted by TiN, and although its strength and hardness is very high, they are slightly inferior to Al2O3. These two aspects led to a degradation of the mechanical properties of the composites, whose strength and hardness depended mainly on their elemental properties. The strength and micro-hardness of samples in group B are shown in Fig. 8. In contrast to group A, the degradation of mechanical performance did not occur in group B, but instead, the bending strength and micro-hardness improved consistent with added TiN. The densities of
Table 4 The relative densities of group B samples. Sample
B0
B1
B2
B3
B4
B5
Density (%)
99.85
99.43
99.92
99.95
99.49
96.96
4
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Fig. 5. SEM (SE) micrographs of the fracture surfaces of TiN–Ti/Al2O3 composites: (a) B0, (b) B4.
four-probe technique in air. A voltage was obtained when the direct current (100 mA) was applied to both ends of the samples using a constant current source. The electrical resistances of the samples were then fitted to the current voltage curve. The electrical conductivities of the samples were calculated with
group B samples imply that the density of sintered samples was not influenced by TiN. Thus, the enhanced mechanical properties of group B samples were primarily attributable to other reasons. The composite properties were key including the reactions among the substrate phases as well as the nature of the reaction products. The emergence of AlTi3N facilitated the combination of the raw-material phases. In addition, a part of Ti was replaced by TiN, which is stronger and harder than Ti. Therefore, the increased TiN contributed to enhanced bending strength and microhardness.
ρ = RS / L = US / LI where L is the length between the two voltage terminals and S is the cross-sectional area of the sample. The results are shown in Figs. 9 and 10. According to the Fig. 9, TiN significantly reduced the resistivity of the composites via its excellent electrical properties. When 2 vol% TiN was added, the resistivity of the composite decreased from
3.4. Effect of TiN on the resistivity of TiN–Ti/Al2O3 composites The electrical resistivities of the samples were measured using a
Fig. 6. EMPA results of the sample B4: (a) all elements, (b) local areas of Ti and N element.
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Fig. 10. The resistivities of group B samples with the increasing TiN content.
Fig. 7. Effects of TiN on the bending strength and micro-hardness of TiN–Ti/Al2O3 composites of group A samples.
resistivity of the composite decreased from an initial value of 3662.4 μΩ cm to a final value of 930.3 μΩ cm with 10 vol% TiN added to a group A sample. The trend in flexural strength and micro-hardness shown in Fig. 7 demonstrate a similar behavior in which their values first increased and then decreased as the amount of TiN increased. The maximum values of the strength and micro-hardness were obtained with 4 vol% and 6 vol% added TiN, respectively. In addition, taking into account the starting point of the cermet, and overall the flexural strength, micro-hardness and resistivities of group A composites, the best effect was achieved with an addition of 5 vol% TiN. As seen in Fig. 10, the conductivities of specimens in group B increased as TiN was added, consistent with those of group A. Compared to group A, the trend was stable when 5 vol% TiN was added, and the resistivities eventually reduced from 2546.5 μΩ cm to 1474.7 μΩ cm. From Figs. 9 and 10, the increased temperature was not conducive to improving the electrical conductivity of composites with the presence of TiN. The strength in group B samples was effectively stable with 15 vol% added TiN, while the micro-hardness did not differ appreciably when the TiN content reached 15 vol% and 20 vol%. Considering the bending strength, micro-hardness, and conductivity, the optimum amount of TiN in group B was 20 vol%. Given these results, composite conductivities were improved by adding TiN due to its excellent electrical conductivity. When TiN was added to composite, TiN particles dispersed to form a conductive network, with the strength of the conductivity depending on the integrity of the conductive network. The trend of resistivity decrease became slow at the later of test, which predicts that the concentration of TiN did not attain its percolation threshold in the composites according to percolation theory. Thus, experiments can be performed in the future using TiN with smaller particles, or changing the Ti and Al2O3 particle sizes.
Fig. 8. Effects of TiN on the bending strength and micro-hardness of TiN–Ti/Al2O3 composites of group B samples.
4. Conclusions Various amount of TiN (as illustrated in Table 1 and 2) were added to Ti/Al2O3 to toughen and strengthen composites by vacuum hotpressing sintering at different temperatures (1400 °C and 1500 °C) under a pressure of 30 MPa. In addition, because a conductive network was generated, the electrical conductivities of composites with added TiN were improved over those of composites without TiN. However, the increased temperature was unfavorable to improving the conductivity because the best resistivity of samples sintered at 1400 °C was 930.3 μΩ cm, while that of a sample sintered at 1500 °C was 1474.7 μΩ cm. Compared with samples without additives, all samples containing TiN had excellent mechanical properties. Optimal results were obtained with 20 vol% added TiN, with a relative density of
Fig. 9. The resistivities of group A samples with the increasing TiN content.
3662.4 μΩ cm to 1749.8 μΩ cm, and the extent of resistivity dropped by half. This agrees well with the samples in group B, which dropped from 2546.5 μΩ cm to 1620.7 μΩ cm as shown in Fig. 10. Subsequently, the resistivities of composites decreased gradually by increasing the amount of TiN added, but this increasing trend decelerated and stabilized. From the overall resistivity trend, the electrical conductivities of the composites was improved remarkably by adding TiN, and the 6
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99.49%, Vickers hardness of 14.94 GPa, flexural strength of 321.55 MPa, and electrical resistivity of 1474.7 μΩ cm. Based on an XRD analysis, the positive effect of TiN may be ascribed to its good performance of high strength and hardness, decreased brittle intermetallic phase, formation of AlTi3N, a cemented carbide, and the finegrained strengthening.
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Acknowledgements This work has been financially supported by the National Natural Science Foundation of China (Grant No. 51372099), the State Key Laboratory Open Fund of Advanced Technology for Materials Synthesis and Processing (Wuhan University of Technology) (Grant No. 2016-KF9) and the Doctoral Found of University of Jinan (XBS1310). References [1] M.A. Tahaa, A.H. Nassara, M.F. Zawrah*, Improvement of wetability, sinterability, mechanical and electrical properties of Al2O3-Ni nanocomposites prepared by mechanical alloying, Ceram. Int. 43 (2017) 3576–3582. [2] L. Rebouta, A. Sousa, P. Capela, M. Andritschky, P. Santilli, A. Matilainen, K. Pischow, N.P. Barradas, E. Alves, Solar selective absorbers based on Al2O3: W cermets and AlSiN/AlSiON layers, Sol. Energy Mater. Sol. Cells 137 (2015) 93–100. [3] M. Zhang, Q.Q. Yang, W.H. Xiong, L.Y. Zheng, B. Huang, S. Chen, Z.H. Yao, Effect of Mo and C additions on magnetic properties of TiC-TiN-Ni cermets, J. Alloy. Compd. 650 (2015) 700–704. [4] H.H. Kusuma, M.K. Saidin, Optical properties of Ti:Al2O3 single crystal, Jab. Fiz. UTM 4 (2009) 42–49. [5] F.Y. Niu, D.J. Wu, S.Y. Zhou, G.Y. Ma, Power prediction for laser engineered net shaping of Al2O3 ceramic parts, J. Eur. Ceram. Soc. 34 (2014) 3811–3817. [6] Y. Li, H.R. Qian, X.D. Cheng, R.F. Zhang, H.P. Zhang, Fabrication of dense cordierite ceramic through reducing Al2O3 mole ratio, Mater. Lett. 116 (2014) 262–264. [7] K. Liua, Y. Shia, C. Lia, L. Haob, J. Liua, Q.S. Wei, Indirect selective laser sintering of epoxy resin- ceramic powers combined with cold isostatic pressing, Ceram. Int. 40 (2014) 7099–7106. [8] R. Muhammad, A. Maurotto, M. Demiral, A. Roy, V. Silberschmidt, Thermally enhanced ultrasonically assisted machining of Ti alloy, CIRP J. Manuf. Sci. Technol. 7 (2014) 159–167. [9] M.J. Liu, Z. Wang, J.Y. Wu, Q.G. Li, Ti/Al2O3 interfacial diffusion: kinetic equation for growth of reaction layer and formation mechanism, J. Alloy. Compd. 652 (2015) 260–265. [10] B.Q. Xie, X.M. Jiang, B.Y. Wang, Z. Cheng, X.G. Peng, Effect of Ti powder addition on strength and oxidation resistance of Al2O3-C materials, Refractories 44 (2010) 34–37.
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Effects of TiN addition on the properties of hot pressed TiN–Ti/Al2O3 composites Qian Liu, Fangfang Qi, Junyan Wu, Liu Zhang, Guopu Shi, Zhi Wang
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School of Material Science and Engineering, University of Jinan, Jinan 250022, China
A R T I C LE I N FO
A B S T R A C T
Keywords: TiN Resistivity Cermet composite Vacuum hot-pressing sintering
TiN–Ti/Al2O3 composites of varying TiN content (0–20 vol%) were prepared by vacuum hot-pressing sintering at different temperatures (1400 °C and 1500 °C) to investigate how TiN affected the mechanical properties and electrical conductivity of the composites. Sintered samples with added TiN exhibited better performance than those without it. The sample with 20 vol% TiN sintered 1500 °C had an optimal relative density of 99.49, Vickers hardness of 14.94 GPa, flexural strength of 321.55 MPa, and electrical resistivity of 1474.7 μΩ cm. However, this increased temperature did not improve the best sample resistivity of 930.3 μΩ cm, which was obtained at 1400 °C. Form SEM images and XRD patterns, the positive effect of TiN on composite mechanical properties may be ascribed to its good performance of high hardness and strength, a decrease of the brittle intermetallic phase, the form of AlTi3N, and the impact of the fine-grained strength of the TiN phase.
1. Introduction Cermet materials have attracted wide attention because their properties originate from both metal and ceramic materials, and include wear and corrosion resistance, high toughness, and strength [1–4]. As one of the most commonly-used high-temperature structural ceramics, Al2O3 is widely used in the semi-conductor, aerospace, and military industries due to its high hardness and good resistance to high temperature, corrosion, and wear [5–7]. On the other hand, metal Ti has been used in the aerospace, marine, and medical fields since it is tough, ductile, and electrically-conductive [8,9]. Most importantly, the two materials are physically and chemically compatible and share similar thermal expansion coefficients. Therefore, they meet the requirements for preparing Ti/Al2O3 cermet, popular in aerodynamicengine and other automotive applications [10–12]. Unfortunately, in long-term studies of Ti/Al2O3 composites, intermetallic compounds of Ti–Al such as TiAl and Ti3Al are always found. These are characterized by low fracture toughness at room temperature and low oxidation resistance at elevated temperatures, significantly restricting the application of Ti/Al2O3 [13–17]. To solve this problem, interface reaction control in Ti–Al materials has been investigated in depth. D. Pilone [13] and Liu et al. [18] reported that Cr, Si, Nb, Y, and W can be alloyed with Ti–Al to reduce the production of Ti–Al compounds by generating other substances. Ai et al. [19] showed that Ti–Al composites could be reinforced by Al2O3 and Nb2O5. Compound Al-TiAl3 materials have also been prepared by
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uniform droplet spraying as described by Wang [14]. Excessive production intermetallic Ti–Al compounds can also be suppressed by adding metals such as Nb or metal oxides such as Y2O3 [20]. For instance, Wu et al. prepared laminar Ti/Al2O3 composites by casting method [21]. In addition, Nb, CeO2, and Y2O3 were doped to reduce the diffusion of atoms at the interface of Ti/Al2O3 composites to restrain the formation of intermetallic Ti–Al compounds on the basis of the lamellar material, by Liu et al. [22,23]. Clearly, despite much effort, the brittleness of intermetallic Ti–Al compounds remains an open issue. Under these circumstances, we find that TiN has long been considered to be ideal in the applications such cutting tools, wear-resistant materials, decorative materials and electrodes owing to its high hardness, beautiful golden color, high chemical stability, and good electrical conductivity [24–26]. Simultaneously, TiN is often used to toughen and strengthen Al2O3 because of its effective pinning. Thus, we hypothesize that when TiN is introduced into Ti/Al2O3 composites, TiN may not only toughen and enhance the composites, but also reduce the generation of Ti–Al intermetallic compounds by combination with N. To date, many studies have centered on TiN, but none has explicitly focused on its effect on TiN–Ti/Al2O3 composites. Therefore, the impact of TiN on the mechanical properties and microstructure of TiN–Ti/ Al2O3 is studied here, and the changes in conductivity and phase with increasing TiN are also characterized.
Corresponding author. E-mail address: [email protected] (Z. Wang).
https://doi.org/10.1016/j.ceramint.2018.03.125 Received 27 February 2018; Received in revised form 13 March 2018; Accepted 14 March 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Liu, Q., Ceramics International (2018), https://doi.org/10.1016/j.ceramint.2018.03.125
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Table 1 Chemical composition of group A samples (wt%). Samples
Ti (vol%)
Al2O3 (vol%)
TiN (vol%)
A0 A1 A2 A3 A4 A5 A6
40 40 40 40 40 40 –
60 58 56 54 52 50 60
– 2 4 6 8 10 40
Table 2 Chemical composition of group B samples (wt%). Samples
Ti (vol%)
Al2O3 (vol%)
TiN (vol%)
B0 B1 B2 B3 B4 B5
40 35 30 25 20 –
60 60 60 60 60 60
– 5 10 15 20 40
Fig. 1. XRD patterns of group A experiment.
2. Experimental procedures In this work, commercially available Al2O3 (99.9% pure, 1.5 µm), Ti (99.9% pure, 20 µm) and TiN (99.5% pure, 1 µm) from ST-NANO C o., Ltd. (Shanghai, China) were used as raw materials. In raw Al2O3 powder, α-Al2O3 accounts for the largest proportion. The experiment was divided into two groups, and the specimen design is summarized in Tables 1 and 2, respectively. Each group includes a TiN/Al2O3 sample in order to explore the reaction between Ti, TiN, and Al2O3. The TiN–Ti/Al2O3 composites were prepared as follows. After weighing, powders were mixed with a planetary ball mill (XQM-2, China) in an ethanol bath (99.7% pure) for 2 h at a milling speed of 200 rpm with a ball-to-powder weight ratio of 2:1. Then, the slurry was dried at 70 °C for 24 h in a vacuum oven. The dry powder was then placed into a graphite mold and sintered by vacuum hot-pressing (VVPgr-80-2300, Shanghai Haoyue, China). Sintering was performed with temperature rise of 10 °C/min (0–1200 °C), and 5 °C/min (1200–1500 °C), at which time the sintering temperature was held at 1400 °C (group A) and 1500 °C (group B) for 1.5 h under a load of 30 MPa. Lastly, sintered samples were removed when the temperature had naturally dropped to about 100 °C, after which the samples were cut, ground, and polished. The relative densities of the samples were determined by the Archimedes method. Using the three-point bending method, the flexural strength of the treated specimens was measured by an electromechanical testing machine (CMT5504, MTS Systems, China) at room temperature with a crosshead speed of 0.5 mm/min and span of 30 mm, with specimens sizes limited to 3 mm × 4 mm × 35 mm. The microhardness was obtained with a Vickers hardness tester (HV-10001S, China) at room temperature with a load of 1000 gf and indentation time was maintained for 15 s. Phase analysis of the samples was performed using an X-ray diffraction analyzer (D8-ADVANCE, Germany) with CuKα radiation, while the microstructure and elemental distributions of the composites were further analyzed with a scanning electron microscope (FEI QUANTA FEG 250, United States), and an electron probe micro-analyzer (Shimadzu EMPA-1600, Japan). In addition, the electrical resistivity was measured using the DC four-probe method with a current-reversal technique [27–29]. All characterization and tests were performed at room temperature.
Fig. 2. XRD patterns of group B experiment.
3. Results and discussion 3.1. Effect of TiN on phase composition of TiN–Ti/Al2O3 composites The results of X-ray diffraction analyses performed to determine the phase constitution of the TiN–Ti/Al2O3 composites are shown in Figs. 1 and 2. The new phase AlTi3N appeared with the addition of TiN, and the diffraction peaks of the TiN phases became stronger with increasing TiN content. No reactions occurred between TiN and Al2O3 for no other phase existed after sintering, as shown at points A6 and B5. As shown in Fig. 1 (A0–A5), the amount of Ti3O gradually reduced with increasing TiN content, because the oxygen diffusion was hindered by the addition of TiN. From A0 to A4, however, the formation of AlTi3N increased gradually, stabilizing at A4 and A5. Meanwhile, the amount of Ti3Al decreased when 2 vol% of TiN was added, as shown in Fig. 1 A1, remaining stable as additional TiN was added. From this, the production of AlTi3N requires two procedures: first, react Ti with Al2O3 to produce Ti3Al, then combine Ti3Al with N to generate AlTi3N. This process also is responsible for reducing the amount of Ti3Al. As shown in Fig. 2, in contrast to samples from group A samples, TiO was found in group B samples sintered at 1500 °C. The amount of TiO decreased until eventually becoming zero as the amount of TiN increased, in line with the trend of Ti3O. The reduction of TiO and Ti3O should be attributed to the decrease of Ti, and the hindering effect of TiN on the diffusion of O elements. In group B, Ti was reduced due to the increased TiN, lowering the Ti3Al. Moreover, a portion of the Ti3Al combined with N to generate AlTi3N. Therefore, the increased TiN led 2
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Fig. 3. SEM (SE) micrographs of the fracture surfaces of group A: A0, (b) A3, (c) A6.
Unfortunately, both caused weak mechanical performance of the composite. As is illustrated in Fig. 3b, when TiN was added to the composite, the phase distribution remained inhomogeneous, while the voids were drastically reduced as compared with the blank sample, and a good combination of phases appeared. From Fig. 3c, the distribution of phases was fairly uniform in the fractured TiN/Al2O3 composite. However, the sample was not visibly densified, because of the requirement for a higher sintering temperature for the TiN/Al2O3 composite. The relative densities of group A samples are listed in Table 3. Overall, the effect of the TiN phase on the density was negligible and the relative density rose slightly because of the reaction between Ti, TiN, and Al2O3. The fracture surfaces of group B samples were observed using SEM in the BSE mode, as shown in Fig. 4. The inhomogeneous phases were still untreated when the sintering temperature reached 1500 °C. However, this gradually improved with the addition of TiN, as shown in Fig. 4b–d, as the increased TiN content led to phase homogenize, especially for the TiN/Al2O3 composite. In addition, the Ti and Al2O3 grain sizes were refined by increasing the TiN content, due to its ability
Table 3 The relative densities of group A samples. Sample
A0
A1
A2
A3
A4
A5
A6
Density (%)
98.32
97.98
98.26
98.96
99.27
99.55
91.59
to a decrease in Ti3Al, but the amount of Ti3Al was too low to be detected when the content of TiN reached 15 vol%. At the same time, the amount of AlTi3N decreased owing to the decline of Ti3Al.
3.2. Effect of TiN on micro-structure of TiN–Ti/Al2O3 composites As shown in Fig. 3, the microstructure of TiN–Ti/Al2O3 cermets with different TiN content sintered at 1400 °C was observed by SEM in the SE mode. Bulk agglomerations are seen in Fig. 3a and massive voids are observed in the enlarged diagram on the right. This implies that Ti and Al2O3 were dispersed very unevenly in the blank sample lacking the TiN phase, while the blank sample A0 had a poor sintering density. 3
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Fig. 4. SEM (BSE) micrographs of the fracture surfaces of group B: (a) B0, (b) B2, (c) B4, (d) B5.
actual test, Ti and N corresponded well, and no separate Ti is shown in Fig. 6. The local Ti and N areas are enlarged and shown in Fig. 6b, indicating that more Ti was present than N despite the overlap between them. Because the amount of Ti was reduced and that of N increased, the local areas marked in the figure are considered to contribute to the TiN phase. Thus, Ti and TiN were mutually attractive during sintering and were mixed together in the sintered sample, reducing the formation of the intermetallic TiAl to some extent.
to pin particles. The bending strength can be enhanced by grain refinement according to the Hall-Petch formula,
σy = σ0 + Kd−1/2 where σy is the yield strength and d is the grain size. In addition, all samples sintered at 1500 °C had high densities with few voids, as shown in Table 4 where the relative densities of all samples were greater than 99%. Thus, the increased sintering temperature promoted the densification of the composites, while the effect of TiN was slight, in good agreement with the change seen in Fig. 4. Comparing Figs. 3c and 4d, the grains in TiN/Al2O3 sintered at 1500 °C were combined better. This confirms the conjecture that the TiN/Al2O3 composite required a higher sintering temperature than Ti/Al2O3. In Fig. 5a, the existence of a crystalline morphology and cracks along grain boundaries are shown, both suggesting that the fracture mode of the composites without added TiN was typically inter-granular. In contrast with B0, the cracks passing through grains in Fig. 5b demonstrate that the fracture mode of B4 with the added 20% TiN had changed from a single inter granular fracture to a mixture of inter granular and transgranular fracture. The change of the fracture mode is indicative of improved composite mechanical properties. In order to observe the distribution of the three main phases, the electron probe micro-analyzer was used to test a section of sample B4, with the results presented in Fig. 6. In Fig. 6a, bright areas are Ti and N element and dark areas are Al and O. In the latter, Al diffused differently from O promoting the formation of AlTi3N. From the XRD pattern, the main phase of the material comprised TiN, Ti, and Al2O3. In the
3.3. Effect of TiN on mechanical properties of TiN–Ti/Al2O3 composites In order to further explore the effects of TiN on Ti/Al2O3 composites, the bending strength and micro-hardness of samples in group A and B were measured, as shown in Figs. 7 and 8 respectively. From Fig. 7, the flexural strength and micro-hardness of group A were consistent. When the TiN content was less than 6 vol%, the decline of density was responsible for the decrease in bending strength and microhardness of samples in the earlier stage. The density increased with added TiN, accompanied by reinforced flexural strength and microhardness. Later, the improved density and weaker mechanical properties appeared simultaneously. This is because early in the addition procedure, the strength and hardness of group A samples depended mainly on their density, but afterwards, the amount of AlTi3N capable of promoting the combination of the three phases (Ti, TiN, Al2O3) did not continue to increase, causing the mechanical performance to degrade. At the same time, Al2O3 was substituted by TiN, and although its strength and hardness is very high, they are slightly inferior to Al2O3. These two aspects led to a degradation of the mechanical properties of the composites, whose strength and hardness depended mainly on their elemental properties. The strength and micro-hardness of samples in group B are shown in Fig. 8. In contrast to group A, the degradation of mechanical performance did not occur in group B, but instead, the bending strength and micro-hardness improved consistent with added TiN. The densities of
Table 4 The relative densities of group B samples. Sample
B0
B1
B2
B3
B4
B5
Density (%)
99.85
99.43
99.92
99.95
99.49
96.96
4
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Fig. 5. SEM (SE) micrographs of the fracture surfaces of TiN–Ti/Al2O3 composites: (a) B0, (b) B4.
four-probe technique in air. A voltage was obtained when the direct current (100 mA) was applied to both ends of the samples using a constant current source. The electrical resistances of the samples were then fitted to the current voltage curve. The electrical conductivities of the samples were calculated with
group B samples imply that the density of sintered samples was not influenced by TiN. Thus, the enhanced mechanical properties of group B samples were primarily attributable to other reasons. The composite properties were key including the reactions among the substrate phases as well as the nature of the reaction products. The emergence of AlTi3N facilitated the combination of the raw-material phases. In addition, a part of Ti was replaced by TiN, which is stronger and harder than Ti. Therefore, the increased TiN contributed to enhanced bending strength and microhardness.
ρ = RS / L = US / LI where L is the length between the two voltage terminals and S is the cross-sectional area of the sample. The results are shown in Figs. 9 and 10. According to the Fig. 9, TiN significantly reduced the resistivity of the composites via its excellent electrical properties. When 2 vol% TiN was added, the resistivity of the composite decreased from
3.4. Effect of TiN on the resistivity of TiN–Ti/Al2O3 composites The electrical resistivities of the samples were measured using a
Fig. 6. EMPA results of the sample B4: (a) all elements, (b) local areas of Ti and N element.
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Fig. 10. The resistivities of group B samples with the increasing TiN content.
Fig. 7. Effects of TiN on the bending strength and micro-hardness of TiN–Ti/Al2O3 composites of group A samples.
resistivity of the composite decreased from an initial value of 3662.4 μΩ cm to a final value of 930.3 μΩ cm with 10 vol% TiN added to a group A sample. The trend in flexural strength and micro-hardness shown in Fig. 7 demonstrate a similar behavior in which their values first increased and then decreased as the amount of TiN increased. The maximum values of the strength and micro-hardness were obtained with 4 vol% and 6 vol% added TiN, respectively. In addition, taking into account the starting point of the cermet, and overall the flexural strength, micro-hardness and resistivities of group A composites, the best effect was achieved with an addition of 5 vol% TiN. As seen in Fig. 10, the conductivities of specimens in group B increased as TiN was added, consistent with those of group A. Compared to group A, the trend was stable when 5 vol% TiN was added, and the resistivities eventually reduced from 2546.5 μΩ cm to 1474.7 μΩ cm. From Figs. 9 and 10, the increased temperature was not conducive to improving the electrical conductivity of composites with the presence of TiN. The strength in group B samples was effectively stable with 15 vol% added TiN, while the micro-hardness did not differ appreciably when the TiN content reached 15 vol% and 20 vol%. Considering the bending strength, micro-hardness, and conductivity, the optimum amount of TiN in group B was 20 vol%. Given these results, composite conductivities were improved by adding TiN due to its excellent electrical conductivity. When TiN was added to composite, TiN particles dispersed to form a conductive network, with the strength of the conductivity depending on the integrity of the conductive network. The trend of resistivity decrease became slow at the later of test, which predicts that the concentration of TiN did not attain its percolation threshold in the composites according to percolation theory. Thus, experiments can be performed in the future using TiN with smaller particles, or changing the Ti and Al2O3 particle sizes.
Fig. 8. Effects of TiN on the bending strength and micro-hardness of TiN–Ti/Al2O3 composites of group B samples.
4. Conclusions Various amount of TiN (as illustrated in Table 1 and 2) were added to Ti/Al2O3 to toughen and strengthen composites by vacuum hotpressing sintering at different temperatures (1400 °C and 1500 °C) under a pressure of 30 MPa. In addition, because a conductive network was generated, the electrical conductivities of composites with added TiN were improved over those of composites without TiN. However, the increased temperature was unfavorable to improving the conductivity because the best resistivity of samples sintered at 1400 °C was 930.3 μΩ cm, while that of a sample sintered at 1500 °C was 1474.7 μΩ cm. Compared with samples without additives, all samples containing TiN had excellent mechanical properties. Optimal results were obtained with 20 vol% added TiN, with a relative density of
Fig. 9. The resistivities of group A samples with the increasing TiN content.
3662.4 μΩ cm to 1749.8 μΩ cm, and the extent of resistivity dropped by half. This agrees well with the samples in group B, which dropped from 2546.5 μΩ cm to 1620.7 μΩ cm as shown in Fig. 10. Subsequently, the resistivities of composites decreased gradually by increasing the amount of TiN added, but this increasing trend decelerated and stabilized. From the overall resistivity trend, the electrical conductivities of the composites was improved remarkably by adding TiN, and the 6
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99.49%, Vickers hardness of 14.94 GPa, flexural strength of 321.55 MPa, and electrical resistivity of 1474.7 μΩ cm. Based on an XRD analysis, the positive effect of TiN may be ascribed to its good performance of high strength and hardness, decreased brittle intermetallic phase, formation of AlTi3N, a cemented carbide, and the finegrained strengthening.
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