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Formation of titanium nitride by mechanical milling and isothermal annealing of titanium and boron nitride
Journal of Alloys and Compounds 391 (2005) 77–81
Formation of titanium nitride by mechanical milling and isothermal annealing of titanium and boron nitride Z.H. Dinga , B. Yaoa,c,∗ , L.X. Qiua , S.Z. Baia , X.Y. Guoa , Y.F. Xuea , W.R. Wanga , X.D. Zhoua , W.H. Sua,b,c b
a Department of Physics, Jilin University, Changchun 130023, PR China Center for the Condensed-matter Science and Technology, Harbin Institute of Technology, Harbin 150001, PR China c International Center for Materials Physics, Academic Sinica 110015, PR China
Received 27 July 2004; received in revised form 30 August 2004; accepted 30 August 2004 Available online 12 October 2004
Abstract The formation of titanium nitride with rock salt structure (␦-TiNx ) was studied by mechanical milling of a mixture of Ti and hexagonal boron nitride (h-BN) powders and isothermal annealing of the mixture after milling for 70 h under argon atmosphere. The mole ratio of Ti to h-BN is 22:78. In the milling process, an amorphous Ti–N alloy was formed firstly by a diffusion reaction between Ti and BN, and then the amorphous Ti–N alloy transformed into ␦-TiNx driven by mechanical milling. However, in the annealing process, a Ti(N) solid solution was formed firstly by incorporation of N into Ti, and the N content in the Ti(N) increased with increasing annealing temperature. When the N content exceeded the solubility limit of the Ti(N) at some annealing temperature, the Ti(N) decomposed into ␦-TiNx and Ti(N) with the solubility limit. No self-sustaining reaction occurs in the present work. No TiB2 is observed to form in the two processes. The thermodynamic and kinetic mechanisms of formation of the ␦-TiNx are discussed. © 2004 Elsevier B.V. All rights reserved. Keywords: Mechanical milling; ␦-TiNx ; Raman spectrum; h-BN; Ti
1. Introduction Refractory materials such as borides, nitrides, carbides and silicides as well as combinations thereof have gained much attention due to their extraordinary hardness; wear resistance, and stability at very high temperature. Titanium nitride with rock salt crystalline structure, termed as ␦-TiNx (x = 30–50 at.%), is one of such refractory materials. It not only has high melting point, but also exhibits high hardness, excellent wear and corrosion resistances. It is, therefore, used widely as coating for cutting tools and wearing parts. Combining ␦-TiNx with some materials to form composites can improve properties of the materials and even can reveal novel properties. For example, titanium diboride (TiB2 ) is a refractory compound with very high melting ∗
Corresponding author. E-mail address: [email protected] (B. Yao).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.08.051
point, high elastic modulus, and plastic deformation even at high temperature [1]. Therefore, it is believed that combining TiB2 and ␦-TiNx into a composite material might yield a unique combination of high-temperature hardness and stability with adequate ductility/fracture toughness. This composite has potential application for jet engine parts, armor plates and cutting tools, etc. [1]. Boron nitride is good insulating ceramic material with good thermal conductivity and high melting point, but it is poor in hardness, mechanical strength, and oxidation-resistance. Doping ␦-TiNx into BN to form ␦-TiNx /BN composite can reinforce mechanical strength of BN and improve its oxidation-resistance. Due to good conductivity of ␦-TiNx , the ␦-TiNx /BN composite will be a conductor with high melting temperature and its conductivity can be adjusted by controlling the ␦-TiNx contents. Hence the ␦-TiNx /BN composite can be used as a heater for melting various metals or alloys. Furthermore, some novel electrical properties may be obtained in the ␦-
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TiNx /BN composites when ␦-TiNx with nanometer size combines with BN. ␦-TiNx is usually synthesized by self-propagating combustion under high nitrogen pressure and high temperature [2], and activated reactive evaporation [3]. It was reported recently that ␦-TiNx was prepared by mechanical milling of metal Ti in a nitrogen atmosphere [4]. In the above synthesis process, N atoms come from nitrogen gas. In recent years, h-BN solid was used to prepare ␦-TiNx or its composites. It was reported that ␦-TiNx /TiB2 composite was prepared by hot pressure reaction between TiH2 and h-BN at high temperature of 1600 ◦ C and self-propagating reaction between Ti and h-BN at a high temperature of 1200 ◦ C [1,5]. Shim et al. [6] reported that ␦-TiNx /TiB2 nanocomposite powder was prepared by high-energy ball milling of a mixture with a Ti to BN ratio of 3:2 for 2 h, and attributed formation of the ␦-TiNx /TiB2 nanocomposite to a self-propagating reaction between Ti and BN during the milling process. However, Li et al. did not obtain the ␦-TiNx /TiB2 nanocomposite by mechanical milling of the same mixture as Shim used, but only obtained a small amount of ␦-TiNx after milling for 40 h. Theses facts indicate that the formation of the ␦-TiNx and TiB2 is related to the experimental conditions used. Conditions and mechanisms of formation of ␦-TiNx and TiB2 by solid reaction between Ti and BN are still problems and need to be investigated. In the present experiment, evolution and products of reaction between Ti and BN are investigated by means of mechanical milling or isothermal annealing techniques, and the mechanisms of formation of ␦-TiNx and TiB2 are discussed based on thermodynamics and kinetics.
2. Experimental procedures The raw materials used in this study were 99% pure Ti powder and 99% pure hexagonal BN powder. The mixture with an atomic ratio of BN to Ti of 78:22 was mechanically milled under Ar atmosphere in a high-energy ball-mill, using stainless balls and vial. The diameters of the balls are 5–15 mm. The ball-to-powder weight ratio was approximately 15:1. The as-milled mixture was heated isothermally at different temperatures and 10−2 Pa. The structure and grain size of the samples were investigated by using an X-ray diffractometer (XRD) with Cu K␣ radiation, and the Raman spectrum of the ␦-TiNx are recorded by Raman spectroscopy (RS) with a 632.8 nm line of a He–Ne gas laser and output power of 20 mW. Measurement of crystallization temperature was performed by differential scanning calorimetry (DSC) with a heating rate of 20 ◦ C/min in a flowing argon atmosphere.
Fig. 1. XRD patterns of the mixture of Ti and h-BN powder milled for (a) 0 h; (b) 50 h; (c) 70 h; (d) 120 h; (e) 160 h, respectively. () h-BN; (䊉) Ti; () ␦-TiNx .
the mixture consists of Ti and h-BN. After milling for 50 h, as shown in Fig. 1(b), all diffraction peaks of h-BN disappear and a diffuse diffraction peak ranging from about 2θ = 10◦ to 30◦ is observed, indicating that h-BN is transformed into amorphous BN (a-BN). The diffraction peak intensity of the Ti decreases and the full width at half-maximum (FWHM) increases, but the peak positions do not shift. In fact, no shift was observed for the peak positions of Ti until the Ti reacted completely with BN during the milling process, as shown in Fig. 1(b)–(d), implying that no Ti solid solution is formed in the whole milling process. It is noted that there is a weak diffuse diffraction peak in diffraction angles of about 30–50◦ in Fig. 1(b), implying that an amorphous alloy may be formed. In order to confirm the formation of the amorphous alloy, a DSC measurement of the 50 h-milled mixture was carried out. An exothermic peak at a peak temperature of 375 ◦ C was observed in the DSC curve, as shown in Fig. 2. That means formation of an amorphous alloy in the 50 h-milled mixture. The crystallization temperature of the amorphous alloy is 375 ◦ C.
3. Experimental results and discussion Fig. 1 shows XRD patterns of the mixture of Ti and h-BN unmilled and milled for different times. Fig. 1(a) shows that
Fig. 2. DSC curve of the mixture milled for 50 h.
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Fig. 4. Dependences of grain size and lattice constant of the ␦-TiNx on milling time.
Fig. 3. Raman spectra of mixture milled for (a) 50 h; (b) 70 h; (c) 100 h; (d) 140 h; (e) 180 h, respectively.
In order to identify the amorphous alloy, Raman scattering measurement was performed for the 50 h-milled mixture, and its Raman spectrum is shown in Fig. 3(a). The spectrum is similar to that of TiN [7], implying that the amorphous alloy is Ti–N amorphous alloy and the Ti–N amorphous alloy has similar Raman spectrum as crystalline ␦-TiN. Extending the milling time to 70 h, as shown in Fig. 1(c), the XRD intensity of the amorphous Ti–N alloy increases, implying more amorphous Ti–N alloy is formed. In addition, a weak diffraction peak is observed at 2θ = 42.1◦ , which almost is the same as the position of the strongest diffraction peak (2 0 0) of ␦-TiNx . Fig. 3(b) shows spectrum of the 70 hmilled mixture, which is similar to Fig. 3(a), but the widths of the peaks are narrower than those in Fig. 3(a). Therefore, it is believed that ␦-TiNx began to form when milling for 70 h. It is due to crystallization of the amorphous Ti–N alloy driven by mechanical milling [8]. With increasing the milling time, the diffraction peak intensities of the Ti become weaker and its peak width is broader, while diffraction peak intensities of the ␦-TiNx increase and the peaks narrow, as shown in Fig. 1(d), implying that more Ti reacts with N of the a-BN to form ␦-TiNx and the grain size of the ␦-TiNx increases as milling time rises, as shown in Fig. 4. The grain size was calculated by Scherrer formula: D = 0.89λ/B cosθ, where λ is the X-ray wavelength, B the FWHM and θ the Bragg diffraction angle. Fig. 3 shows that the width of Raman scattering peaks
decrease with increasing grain size, similar to XRD results in Fig. 1. It is also found from the XRD results that there always exists the diffraction peak of the amorphous Ti–N alloy for samples prepared at milling times less than 140 h, where Ti does not react completely with a-BN. Upon milling for 160 h, the diffraction peaks of both Ti and the amorphous Ti–N alloy disappeared completely. The XRD pattern consists of diffraction peaks of a-BN and ␦-TiNx , as shown in Fig. 1(e). That indicates that all Ti reacts with N in the a-BN to form ␦-TiNx . The lattice constant of the ␦-TiNx produced at different milling times was determined, and is shown in Fig. 4, indicating that the lattice constant decreases with increasing milling time. ␦-TiNx has rock salt structure and exists in a N content range of about 30–50 at.%. For Ti-rich ␦-TiNx the excess Ti occupy N positions. Since the atomic radius of N is less than that of Ti, the Ti-rich ␦-TiNx has larger lattice constant, and the lattice constant decreases with increasing N content. Therefore, the decrease of the lattice constant of ␦-TiNx in Fig. 4 is attributed to an increment of the N content in ␦-TiNx with increasing milling time. Based on the discussion mentioned above, it is deduced that ␦-TiNx is formed by a diffuse reaction between Ti and a-BN and a phase transition from amorphous Ti–N alloy to ␦-TiNx in ball milling process. Firstly, the N atoms of a-BN become incorporated into the nanocrystalline Ti with a large amount of defects induced during the milling process to form amorphous Ti–N alloy. It is estimated that the local temperature induced by the collisions of the steel balls is in the range of 420–550 ◦ C and the local pressure is 3–4 GPa [8]. The local temperature is higher than the crystallization temperature of the amorphous Ti–N alloy, 375 ◦ C, while the local pressure can promote crystallization of the amorphous alloy in a polymorphous crystallization mode. Therefore, when the N content in the amorphous alloy exceeds some critical value, which should be above 30 at.% [9], the amorphous Ti–N crystallizes to ␦-TiNx driven by the local temperature and local pressure [10]. With extended milling time, on the one hand,
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the N atoms continue to dissolve into the nanocrystalline Ti to form amorphous Ti–N alloy. On the other hand, the N atoms also become incorporated into the amorphous Ti–N and into ␦-TiNx , leading to an increment of the N content of ␦-TiNx and hence to a decrease of its lattice constant with increasing milling time. However, no evidence shows that ␦TiNx is formed by a self-sustaining reaction in the present experiment. In order to study the mechanism of formation of ␦-TiNx by a solid reaction between Ti and a-BN, the 70 h-milled mixture, which is mainly composed of Ti and a-BN besides a small amount of amorphous Ti–N and ␦-TiNx (Fig. 1(c)), was annealed for 1 h under 10−2 Pa in the temperature ranging from 470 to 950 ◦ C. As shown in Fig. 5, the XRD peaks of ␦-TiNx are observed more clearly and the XRD intensities become stronger with increasing annealing temperature. On the other hand, all diffraction peaks of Ti broaden asymmetrically for the samples prepared at annealing temperatures of 470–550 ◦ C, as shown in Fig. 5(a)–(c). For each of the peaks, the diffraction intensity decreases towards lower diffraction angles, resulting in the diffraction peaks of Ti becoming asymmetric. This asymmetric behavior is attributed to the N atoms from a-BN that is incorporated inhomogeneously into Ti to form Ti(N) solid solutions with different lattice constants. This is obvious for increasing annealing temperatures up to 550 ◦ C, indicating that more Ti(N) solid
Fig. 5. XRD patterns of 70 h-milled mixture annealed at (a) 470 ◦ C; (b) 500 ◦ C; (c) 550 ◦ C; (d) 600 ◦ C; (e) 700 ◦ C; (f) 800 ◦ C; (g) 950 ◦ C, respectively. (䊉) Ti; () ␦-TiNx .
solutions are formed and the N content in the Ti(N) increases with increasing annealing temperature. However, when annealed above 600 ◦ C, the asymmetrical behavior disappears, and the diffraction peak positions of Ti(N) shift toward lower diffraction angle, as shown in Fig. 5(d), then retain their positions for annealing temperatures of 600–950 ◦ C, as shown in Fig. 5(d)–(g). The diffraction peak near 37◦ can be split into two peaks, the one is located at 37.39◦ , belongs to the (0 0 2) reflection of Ti(N), and the other peak corresponds to (1 1 1) reflection of ␦-TiNx . However, its peak position increases from 36.21◦ for the sample produced at an annealing temperature of 600 ◦ C to 36.58◦ for the sample prepared at 950 ◦ C. That implies the Ti(N) solid solution begins to decompose into the Ti(N) and ␦-TiNx phases when the annealing temperature exceeds 600 ◦ C, because the N content in the Ti(N) exceeds its solubility limit, which is about 12 at.% at 600 ◦ C. The decrease of the lattice constant of ␦-TiNx with annealing temperature is due to the increment of the N content in ␦-TiNx , as the atomic radius of N is smaller than that of Ti. Based on discussion mentioned above, the formation of ␦-TiNx originates from two processes in the isothermally annealing process. One is the crystallization of the amorphous Ti–N at the annealing temperatures of 470–550 ◦ C. In this process only a small amount of ␦-TiNx is produced. Another is the decomposition of the Ti(N) solid solution upon annealing above 600 ◦ C, where most of the ␦-TiNx is formed. The mechanism of formation of the ␦-TiNx in the isothermally annealing process is obviously different from that in the mechanical milling process. It was reported that TiB2 and ␦-TiNx were prepared by mechanical milling of a mixture with a mole ratio of Ti to BN of 3:2 or by annealing of the mixture at high temperatures of above 1200 ◦ C [1,5]. However, no TiB2 was observed to form in the present experiment during mechanical milling of the mixture of Ti and BN with various mole ratios or annealing of the mixture at temperatures up to 950 ◦ C. It is well known that heat of formation of TiN is −81 kcal/mol, which is more negative than the heat of formation of TiB2 of −67 kcal/mol [11]. So from the point of view of thermodynamics, TiN should form more favorably than TiB2 when Ti reacts with a-BN. On the other hand, it is found from the Ti–B and Ti–N phase diagrams [9] that N atoms can become incorporated into Ti to form a Ti(N) solid solution but the B atoms cannot. Hence, the N atoms react easier with Ti to form amorphous Ti–N alloy or a Ti(N) solid solution than B atoms to form amorphous Ti–B alloy from the point of view of kinetics. Therefore, it is favorable thermodynamically and kinetically for Ti to react with N from the a-BN to form ␦TiNx during mechanical milling or annealing of the mixture of Ti and BN. A similar result was found in the preparation of Fe–N alloy by mechanical milling or annealing of the mixture of Fe and BN [12]. During milling or annealing some N atoms from the a-BN react with Ti to form ␦-TiNx with the result that the remaining a-BN becomes B-rich amorphous BN. That was confirmed by DSC measurement. It was found that the crystallization
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temperature at atmospheric pressure is about 630 ◦ C for aBN, but above 900 ◦ C for the B-rich amorphous BN. This fact shows that the remaining a-BN is not a-BN but B-rich amorphous BN.
the annealing process. No TiB2 was observed to form during milling or annealing. That is due to the fact that the heat of formation of TiN is more negative than that of TiB2 and the solubility limit is zero for B in Ti but is not for N in Ti.
4. Conclusions
References
Nanocrystalline ␦-TiNx was prepared by mechanical milling of a mixture of Ti and BN or annealing of this mixture after milling for 70 h. The ␦-TiNx is formed by a diffusive reaction between Ti and a-BN and a phase transition but not by a self-propagating reaction during the milling process in the present experiment. Firstly, Ti reacts with N in a-BN to form an amorphous Ti–N alloy, and then the latter crystallizes ␦-TiNx driven by the local pressure and local temperature induced by the collisions between the balls or the ball and the vial. The formation of ␦-TiNx during annealing is attributed to two processes. Below 600 ◦ C, a small amount of ␦-TiNx is formed by crystallization of the amorphous Ti–N alloy in the 70 h-milled mixture. When annealed above 600 ◦ C, the Ti(N) solid solutions with N content higher than the solubility limit decompose into a Ti(N) solid solution corresponding to the solubility limit and ␦-TiNx . Most of the ␦-TiNx compound is formed by decomposition of the Ti(N) solid solution in
[1] F. Olevsky, P. Mogilevsky, E.Y. Gutmanas, I. Gotman, Metall. Mater. Trans. A 27A (1996) 2071. [2] S. Deevi, Z.A. Munir, J. Mater. Res. 5 (1990) 2177. [3] K. Nakamura, K. Inagawa, K. Tsuruoka, S. Komiya, Thin Solid Films 40 (1997) 155. [4] Z.H. Chin, T.P. Perng, Appl. Phys. Lett. 70 (18) (1997) 2380. [5] G.J. Zhang, J. Am. Ceram. Soc. 67 (2) (1988) 342. [6] J.H. Shim, J.S. Byun, Y.W. Cho, Scripta Mater. 47 (2002) 493. [7] Y.H. Cheng, B.K. Tay, S.P. Lau, H. Kuper, F. Richter, J. Appl. Phys. 92 (4) (2002) 1845. [8] B. Yao, S.E. Liu, L. Liu, L. Si, W.H. Su, Y. Li, J. Appl. Phys. 90 (3) (2001) 1650. [9] J.Q. Yu, W.Z. Yi, B.D. Chen, H.J. Chen, Phase Diagram of Binary Alloys, Science and Technology Press of Shanghai, Shanghai, 1987, p. 479. [10] B. Yao, L. Liu, S.E. Liu, W.H. Su, Y. Li, J. Non-cryst. Solids 277 (2/3) (2000) 91–97. [11] I. Bain, O. Knacke, Thermochemical Properties of Inorganic Substances, Springer-Verlag, New York, 1973, p. 674, 749. [12] B. Yao, L. Liu, S.E. Liu, X. Hu, W.H. Su, J. Mater. Res. 16 (9) (2001) 2583.
Formation of titanium nitride by mechanical milling and isothermal annealing of titanium and boron nitride Z.H. Dinga , B. Yaoa,c,∗ , L.X. Qiua , S.Z. Baia , X.Y. Guoa , Y.F. Xuea , W.R. Wanga , X.D. Zhoua , W.H. Sua,b,c b
a Department of Physics, Jilin University, Changchun 130023, PR China Center for the Condensed-matter Science and Technology, Harbin Institute of Technology, Harbin 150001, PR China c International Center for Materials Physics, Academic Sinica 110015, PR China
Received 27 July 2004; received in revised form 30 August 2004; accepted 30 August 2004 Available online 12 October 2004
Abstract The formation of titanium nitride with rock salt structure (␦-TiNx ) was studied by mechanical milling of a mixture of Ti and hexagonal boron nitride (h-BN) powders and isothermal annealing of the mixture after milling for 70 h under argon atmosphere. The mole ratio of Ti to h-BN is 22:78. In the milling process, an amorphous Ti–N alloy was formed firstly by a diffusion reaction between Ti and BN, and then the amorphous Ti–N alloy transformed into ␦-TiNx driven by mechanical milling. However, in the annealing process, a Ti(N) solid solution was formed firstly by incorporation of N into Ti, and the N content in the Ti(N) increased with increasing annealing temperature. When the N content exceeded the solubility limit of the Ti(N) at some annealing temperature, the Ti(N) decomposed into ␦-TiNx and Ti(N) with the solubility limit. No self-sustaining reaction occurs in the present work. No TiB2 is observed to form in the two processes. The thermodynamic and kinetic mechanisms of formation of the ␦-TiNx are discussed. © 2004 Elsevier B.V. All rights reserved. Keywords: Mechanical milling; ␦-TiNx ; Raman spectrum; h-BN; Ti
1. Introduction Refractory materials such as borides, nitrides, carbides and silicides as well as combinations thereof have gained much attention due to their extraordinary hardness; wear resistance, and stability at very high temperature. Titanium nitride with rock salt crystalline structure, termed as ␦-TiNx (x = 30–50 at.%), is one of such refractory materials. It not only has high melting point, but also exhibits high hardness, excellent wear and corrosion resistances. It is, therefore, used widely as coating for cutting tools and wearing parts. Combining ␦-TiNx with some materials to form composites can improve properties of the materials and even can reveal novel properties. For example, titanium diboride (TiB2 ) is a refractory compound with very high melting ∗
Corresponding author. E-mail address: [email protected] (B. Yao).
0925-8388/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2004.08.051
point, high elastic modulus, and plastic deformation even at high temperature [1]. Therefore, it is believed that combining TiB2 and ␦-TiNx into a composite material might yield a unique combination of high-temperature hardness and stability with adequate ductility/fracture toughness. This composite has potential application for jet engine parts, armor plates and cutting tools, etc. [1]. Boron nitride is good insulating ceramic material with good thermal conductivity and high melting point, but it is poor in hardness, mechanical strength, and oxidation-resistance. Doping ␦-TiNx into BN to form ␦-TiNx /BN composite can reinforce mechanical strength of BN and improve its oxidation-resistance. Due to good conductivity of ␦-TiNx , the ␦-TiNx /BN composite will be a conductor with high melting temperature and its conductivity can be adjusted by controlling the ␦-TiNx contents. Hence the ␦-TiNx /BN composite can be used as a heater for melting various metals or alloys. Furthermore, some novel electrical properties may be obtained in the ␦-
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TiNx /BN composites when ␦-TiNx with nanometer size combines with BN. ␦-TiNx is usually synthesized by self-propagating combustion under high nitrogen pressure and high temperature [2], and activated reactive evaporation [3]. It was reported recently that ␦-TiNx was prepared by mechanical milling of metal Ti in a nitrogen atmosphere [4]. In the above synthesis process, N atoms come from nitrogen gas. In recent years, h-BN solid was used to prepare ␦-TiNx or its composites. It was reported that ␦-TiNx /TiB2 composite was prepared by hot pressure reaction between TiH2 and h-BN at high temperature of 1600 ◦ C and self-propagating reaction between Ti and h-BN at a high temperature of 1200 ◦ C [1,5]. Shim et al. [6] reported that ␦-TiNx /TiB2 nanocomposite powder was prepared by high-energy ball milling of a mixture with a Ti to BN ratio of 3:2 for 2 h, and attributed formation of the ␦-TiNx /TiB2 nanocomposite to a self-propagating reaction between Ti and BN during the milling process. However, Li et al. did not obtain the ␦-TiNx /TiB2 nanocomposite by mechanical milling of the same mixture as Shim used, but only obtained a small amount of ␦-TiNx after milling for 40 h. Theses facts indicate that the formation of the ␦-TiNx and TiB2 is related to the experimental conditions used. Conditions and mechanisms of formation of ␦-TiNx and TiB2 by solid reaction between Ti and BN are still problems and need to be investigated. In the present experiment, evolution and products of reaction between Ti and BN are investigated by means of mechanical milling or isothermal annealing techniques, and the mechanisms of formation of ␦-TiNx and TiB2 are discussed based on thermodynamics and kinetics.
2. Experimental procedures The raw materials used in this study were 99% pure Ti powder and 99% pure hexagonal BN powder. The mixture with an atomic ratio of BN to Ti of 78:22 was mechanically milled under Ar atmosphere in a high-energy ball-mill, using stainless balls and vial. The diameters of the balls are 5–15 mm. The ball-to-powder weight ratio was approximately 15:1. The as-milled mixture was heated isothermally at different temperatures and 10−2 Pa. The structure and grain size of the samples were investigated by using an X-ray diffractometer (XRD) with Cu K␣ radiation, and the Raman spectrum of the ␦-TiNx are recorded by Raman spectroscopy (RS) with a 632.8 nm line of a He–Ne gas laser and output power of 20 mW. Measurement of crystallization temperature was performed by differential scanning calorimetry (DSC) with a heating rate of 20 ◦ C/min in a flowing argon atmosphere.
Fig. 1. XRD patterns of the mixture of Ti and h-BN powder milled for (a) 0 h; (b) 50 h; (c) 70 h; (d) 120 h; (e) 160 h, respectively. () h-BN; (䊉) Ti; () ␦-TiNx .
the mixture consists of Ti and h-BN. After milling for 50 h, as shown in Fig. 1(b), all diffraction peaks of h-BN disappear and a diffuse diffraction peak ranging from about 2θ = 10◦ to 30◦ is observed, indicating that h-BN is transformed into amorphous BN (a-BN). The diffraction peak intensity of the Ti decreases and the full width at half-maximum (FWHM) increases, but the peak positions do not shift. In fact, no shift was observed for the peak positions of Ti until the Ti reacted completely with BN during the milling process, as shown in Fig. 1(b)–(d), implying that no Ti solid solution is formed in the whole milling process. It is noted that there is a weak diffuse diffraction peak in diffraction angles of about 30–50◦ in Fig. 1(b), implying that an amorphous alloy may be formed. In order to confirm the formation of the amorphous alloy, a DSC measurement of the 50 h-milled mixture was carried out. An exothermic peak at a peak temperature of 375 ◦ C was observed in the DSC curve, as shown in Fig. 2. That means formation of an amorphous alloy in the 50 h-milled mixture. The crystallization temperature of the amorphous alloy is 375 ◦ C.
3. Experimental results and discussion Fig. 1 shows XRD patterns of the mixture of Ti and h-BN unmilled and milled for different times. Fig. 1(a) shows that
Fig. 2. DSC curve of the mixture milled for 50 h.
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Fig. 4. Dependences of grain size and lattice constant of the ␦-TiNx on milling time.
Fig. 3. Raman spectra of mixture milled for (a) 50 h; (b) 70 h; (c) 100 h; (d) 140 h; (e) 180 h, respectively.
In order to identify the amorphous alloy, Raman scattering measurement was performed for the 50 h-milled mixture, and its Raman spectrum is shown in Fig. 3(a). The spectrum is similar to that of TiN [7], implying that the amorphous alloy is Ti–N amorphous alloy and the Ti–N amorphous alloy has similar Raman spectrum as crystalline ␦-TiN. Extending the milling time to 70 h, as shown in Fig. 1(c), the XRD intensity of the amorphous Ti–N alloy increases, implying more amorphous Ti–N alloy is formed. In addition, a weak diffraction peak is observed at 2θ = 42.1◦ , which almost is the same as the position of the strongest diffraction peak (2 0 0) of ␦-TiNx . Fig. 3(b) shows spectrum of the 70 hmilled mixture, which is similar to Fig. 3(a), but the widths of the peaks are narrower than those in Fig. 3(a). Therefore, it is believed that ␦-TiNx began to form when milling for 70 h. It is due to crystallization of the amorphous Ti–N alloy driven by mechanical milling [8]. With increasing the milling time, the diffraction peak intensities of the Ti become weaker and its peak width is broader, while diffraction peak intensities of the ␦-TiNx increase and the peaks narrow, as shown in Fig. 1(d), implying that more Ti reacts with N of the a-BN to form ␦-TiNx and the grain size of the ␦-TiNx increases as milling time rises, as shown in Fig. 4. The grain size was calculated by Scherrer formula: D = 0.89λ/B cosθ, where λ is the X-ray wavelength, B the FWHM and θ the Bragg diffraction angle. Fig. 3 shows that the width of Raman scattering peaks
decrease with increasing grain size, similar to XRD results in Fig. 1. It is also found from the XRD results that there always exists the diffraction peak of the amorphous Ti–N alloy for samples prepared at milling times less than 140 h, where Ti does not react completely with a-BN. Upon milling for 160 h, the diffraction peaks of both Ti and the amorphous Ti–N alloy disappeared completely. The XRD pattern consists of diffraction peaks of a-BN and ␦-TiNx , as shown in Fig. 1(e). That indicates that all Ti reacts with N in the a-BN to form ␦-TiNx . The lattice constant of the ␦-TiNx produced at different milling times was determined, and is shown in Fig. 4, indicating that the lattice constant decreases with increasing milling time. ␦-TiNx has rock salt structure and exists in a N content range of about 30–50 at.%. For Ti-rich ␦-TiNx the excess Ti occupy N positions. Since the atomic radius of N is less than that of Ti, the Ti-rich ␦-TiNx has larger lattice constant, and the lattice constant decreases with increasing N content. Therefore, the decrease of the lattice constant of ␦-TiNx in Fig. 4 is attributed to an increment of the N content in ␦-TiNx with increasing milling time. Based on the discussion mentioned above, it is deduced that ␦-TiNx is formed by a diffuse reaction between Ti and a-BN and a phase transition from amorphous Ti–N alloy to ␦-TiNx in ball milling process. Firstly, the N atoms of a-BN become incorporated into the nanocrystalline Ti with a large amount of defects induced during the milling process to form amorphous Ti–N alloy. It is estimated that the local temperature induced by the collisions of the steel balls is in the range of 420–550 ◦ C and the local pressure is 3–4 GPa [8]. The local temperature is higher than the crystallization temperature of the amorphous Ti–N alloy, 375 ◦ C, while the local pressure can promote crystallization of the amorphous alloy in a polymorphous crystallization mode. Therefore, when the N content in the amorphous alloy exceeds some critical value, which should be above 30 at.% [9], the amorphous Ti–N crystallizes to ␦-TiNx driven by the local temperature and local pressure [10]. With extended milling time, on the one hand,
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the N atoms continue to dissolve into the nanocrystalline Ti to form amorphous Ti–N alloy. On the other hand, the N atoms also become incorporated into the amorphous Ti–N and into ␦-TiNx , leading to an increment of the N content of ␦-TiNx and hence to a decrease of its lattice constant with increasing milling time. However, no evidence shows that ␦TiNx is formed by a self-sustaining reaction in the present experiment. In order to study the mechanism of formation of ␦-TiNx by a solid reaction between Ti and a-BN, the 70 h-milled mixture, which is mainly composed of Ti and a-BN besides a small amount of amorphous Ti–N and ␦-TiNx (Fig. 1(c)), was annealed for 1 h under 10−2 Pa in the temperature ranging from 470 to 950 ◦ C. As shown in Fig. 5, the XRD peaks of ␦-TiNx are observed more clearly and the XRD intensities become stronger with increasing annealing temperature. On the other hand, all diffraction peaks of Ti broaden asymmetrically for the samples prepared at annealing temperatures of 470–550 ◦ C, as shown in Fig. 5(a)–(c). For each of the peaks, the diffraction intensity decreases towards lower diffraction angles, resulting in the diffraction peaks of Ti becoming asymmetric. This asymmetric behavior is attributed to the N atoms from a-BN that is incorporated inhomogeneously into Ti to form Ti(N) solid solutions with different lattice constants. This is obvious for increasing annealing temperatures up to 550 ◦ C, indicating that more Ti(N) solid
Fig. 5. XRD patterns of 70 h-milled mixture annealed at (a) 470 ◦ C; (b) 500 ◦ C; (c) 550 ◦ C; (d) 600 ◦ C; (e) 700 ◦ C; (f) 800 ◦ C; (g) 950 ◦ C, respectively. (䊉) Ti; () ␦-TiNx .
solutions are formed and the N content in the Ti(N) increases with increasing annealing temperature. However, when annealed above 600 ◦ C, the asymmetrical behavior disappears, and the diffraction peak positions of Ti(N) shift toward lower diffraction angle, as shown in Fig. 5(d), then retain their positions for annealing temperatures of 600–950 ◦ C, as shown in Fig. 5(d)–(g). The diffraction peak near 37◦ can be split into two peaks, the one is located at 37.39◦ , belongs to the (0 0 2) reflection of Ti(N), and the other peak corresponds to (1 1 1) reflection of ␦-TiNx . However, its peak position increases from 36.21◦ for the sample produced at an annealing temperature of 600 ◦ C to 36.58◦ for the sample prepared at 950 ◦ C. That implies the Ti(N) solid solution begins to decompose into the Ti(N) and ␦-TiNx phases when the annealing temperature exceeds 600 ◦ C, because the N content in the Ti(N) exceeds its solubility limit, which is about 12 at.% at 600 ◦ C. The decrease of the lattice constant of ␦-TiNx with annealing temperature is due to the increment of the N content in ␦-TiNx , as the atomic radius of N is smaller than that of Ti. Based on discussion mentioned above, the formation of ␦-TiNx originates from two processes in the isothermally annealing process. One is the crystallization of the amorphous Ti–N at the annealing temperatures of 470–550 ◦ C. In this process only a small amount of ␦-TiNx is produced. Another is the decomposition of the Ti(N) solid solution upon annealing above 600 ◦ C, where most of the ␦-TiNx is formed. The mechanism of formation of the ␦-TiNx in the isothermally annealing process is obviously different from that in the mechanical milling process. It was reported that TiB2 and ␦-TiNx were prepared by mechanical milling of a mixture with a mole ratio of Ti to BN of 3:2 or by annealing of the mixture at high temperatures of above 1200 ◦ C [1,5]. However, no TiB2 was observed to form in the present experiment during mechanical milling of the mixture of Ti and BN with various mole ratios or annealing of the mixture at temperatures up to 950 ◦ C. It is well known that heat of formation of TiN is −81 kcal/mol, which is more negative than the heat of formation of TiB2 of −67 kcal/mol [11]. So from the point of view of thermodynamics, TiN should form more favorably than TiB2 when Ti reacts with a-BN. On the other hand, it is found from the Ti–B and Ti–N phase diagrams [9] that N atoms can become incorporated into Ti to form a Ti(N) solid solution but the B atoms cannot. Hence, the N atoms react easier with Ti to form amorphous Ti–N alloy or a Ti(N) solid solution than B atoms to form amorphous Ti–B alloy from the point of view of kinetics. Therefore, it is favorable thermodynamically and kinetically for Ti to react with N from the a-BN to form ␦TiNx during mechanical milling or annealing of the mixture of Ti and BN. A similar result was found in the preparation of Fe–N alloy by mechanical milling or annealing of the mixture of Fe and BN [12]. During milling or annealing some N atoms from the a-BN react with Ti to form ␦-TiNx with the result that the remaining a-BN becomes B-rich amorphous BN. That was confirmed by DSC measurement. It was found that the crystallization
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temperature at atmospheric pressure is about 630 ◦ C for aBN, but above 900 ◦ C for the B-rich amorphous BN. This fact shows that the remaining a-BN is not a-BN but B-rich amorphous BN.
the annealing process. No TiB2 was observed to form during milling or annealing. That is due to the fact that the heat of formation of TiN is more negative than that of TiB2 and the solubility limit is zero for B in Ti but is not for N in Ti.
4. Conclusions
References
Nanocrystalline ␦-TiNx was prepared by mechanical milling of a mixture of Ti and BN or annealing of this mixture after milling for 70 h. The ␦-TiNx is formed by a diffusive reaction between Ti and a-BN and a phase transition but not by a self-propagating reaction during the milling process in the present experiment. Firstly, Ti reacts with N in a-BN to form an amorphous Ti–N alloy, and then the latter crystallizes ␦-TiNx driven by the local pressure and local temperature induced by the collisions between the balls or the ball and the vial. The formation of ␦-TiNx during annealing is attributed to two processes. Below 600 ◦ C, a small amount of ␦-TiNx is formed by crystallization of the amorphous Ti–N alloy in the 70 h-milled mixture. When annealed above 600 ◦ C, the Ti(N) solid solutions with N content higher than the solubility limit decompose into a Ti(N) solid solution corresponding to the solubility limit and ␦-TiNx . Most of the ␦-TiNx compound is formed by decomposition of the Ti(N) solid solution in
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