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Deposition of Ti–Al–N coatings by thermal CVD
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International Journal of Refractory Metals & Hard Materials 26 (2008) 563–568 www.elsevier.com/locate/IJRMHM
Deposition of Ti–Al–N coatings by thermal CVD J. Wagner a, V. Edlmayr a, M. Penoy b, C. Michotte b, C. Mitterer c,*, M. Kathrein d a
Materials Center Leoben Forschung GmbH, Roseggerstrasse 12, A-8700 Leoben, Austria b CERATIZIT Luxembourg S.a`.r.l., L-8201 Mamer, Luxembourg c Department of Physical Metallurgy and Materials Testing, University of Leoben, Franz-Josef-Strasse 18, A-8700 Leoben, Austria d CERATIZIT Austria GmbH, A-6600 Reutte, Austria Received 13 December 2007; accepted 17 January 2008
Abstract This study investigates the feasibility to deposit metastable Ti1–xAlxN coatings by thermal chemical vapour deposition (CVD) from the TiCl4–AlCl3–NH3–Ar system in an industrial CVD unit. The chemical composition and phase constitution was altered using different AlCl3 contents and deposition temperatures. According to X-ray diffraction analysis, single-phase face-centered cubic (fcc) structures have been obtained up to 19 at.% Al, where the incorporation of Al decreases grain size and micro-hardness. In addition, fcc-AlN has been detected between 20 and 24 at.% Al causing an enhanced micro-hardness and extremely fine structures. High temperatures and Al-contents produce coatings of poor mechanical properties consisting of TiN and hexagonal close-packed AlN. Ó 2008 Elsevier Ltd. All rights reserved. Keywords: Ti–Al–N; CVD; Structure; Hardness; Aluminium nitride
1. Introduction The deposition of Ti1–xAlxN (0 6 x 6 1) coatings by plasma-assisted chemical vapour deposition (CVD) [1] and physical vapour deposition (PVD) [2] techniques is well-established in cutting tool industry. High hardness and oxidation resistance enable the application in severe machining processes like high-speed and dry cutting [3]. The excellent properties result from the substitution of Ti by Al atoms in the face-centered cubic (fcc) TiN lattice, which improves both mechanical properties and thermal stability of TiN [4]. Although the mutual solubility of fcc-TiN and hexagonal close-packed (hcp) AlN is negligible in thermodynamic equilibrium, the non-equilibrium nature of plasma-assisted processes causes an extended solubility and the formation of metastable phases [5]. From the resulting supersaturated phases, fcc-Ti1–xAlxN (x < 0.6– 0.7) is ideal for cutting applications, while the presence of
*
Corresponding author. E-mail address: [email protected] (C. Mitterer).
0263-4368/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2008.01.003
hcp-AlxTi1–xN is widely accepted to deteriorate the performance [6]. The deposition of TiN for wear-resistant applications by thermal CVD is a conventional process since many years [7]. However, the incorporation of Al and formation of fcc-Ti1–xAlxN coatings without plasma assistance is challenging [8]. Single-phase fcc-Ti1–xAlxN (x < 0.57) coatings have been obtained by thermal CVD using metal-organic precursors [9,10]. Disadvantageous of metal-organic precursors are their high costs, and the low deposition temperatures may cause an insufficient adhesion. Most frequently used in cutting tool industry are the less expensive metal chlorides, e.g. TiCl4 together with N2 and H2 is a customary precursor system for the deposition of TiN. However, the incorporation of Al into TiN using AlCl3, N2, and H2 is thermodynamically not feasible. The utilization of NH3 as nitrogen source and reducing agent instead of N2 and H2 enables the formation of both TiN and AlN within a comparable temperature range [8]. Most studies consider the deposition of either single-phase coatings of fcc-TiN [11] and hcp-AlN [12], respectively, or fcc-TiN/hcp-AlN composites [13], but only few investigations report on
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fcc-Ti1–xAlxN deposits. Metastable fcc coatings have been deposited using TiCl3 and AlCl3 from in-situ chlorination of TiAl-alloys with Cl2 [14,15]. This process requires sophisticated equipment and bears the risk of preferred evaporation of a constituent. Recently, fcc-Ti1–xAlxN (x < 0.9) coatings have been deposited from TiCl4 and AlCl3 together with NH3 and H2 in a laboratory scale unit [16]. Here, we explore the feasibility to deposit Ti–Al–N coatings by thermal CVD, in particular fcc-Ti1–xAlxN, using the conventional precursors TiCl4 and AlCl3 together with NH3 in an industrial deposition plant. The Ti–Al–N coatings have been grown at different temperatures and AlCl3contents. The coatings have been characterised with respect to composition, structure, morphology, and hardness. 2. Experimental The deposition experiments were performed in an industrial vertical hot-wall CVD system (inner diameter of the reactor 205 mm, height 300 mm, usable for up to 1000 cemented carbide indexable inserts) at atmospheric pressure [17] using TiCl4, AlCl3, NH3, and Ar. AlCl3 was generated by reaction between HCl and aluminium chips. TiCl4 was provided by passing Ar through liquid TiCl4. The purity of the employed gases and reactants was better than 99.999%. The deposition unit was evacuated to remove residual gases and leak tested prior to all deposition processes. During the heating step, the deposition chamber was rinsed with H2 to remove organic contaminations and oxide layers on the substrates. Residual reactants and byproducts were removed at the end of the deposition process by rinsing the chamber with H2. The coating composition was adjusted by altering the deposition temperature and the AlCl3 partial pressure p(AlCl3). At a substrate temperature of 550 °C, p(AlCl3) was varied between 0 and 246 Pa, while other parameters were kept constant. The experiments between 600 and 700 °C were performed with half the concentration of reactive gases (i.e., AlCl3, TiCl4, and NH3) for improved growth rate homogeneity. Thus, p(AlCl3) was altered between 0 and 123 Pa at 600 °C. For the depositions at 650 and 700 °C, p(AlCl3) was set to 62 Pa. The coating deposition was performed on cemented carbide specimens (77 wt.% WC, 11 wt.% Co, 4 wt.% TiC, 8 wt.% (Ta,Nb)C). The sample position was the same throughout all deposition runs. The deposition time was 3 h for deposition temperatures Tdep of 550 and 600 °C and 2 h at Tdep of 650 and 700 °C. The mean chemical composition of the coatings was measured from glow discharge optical emission spectroscopy (GDOES) depth-profiles (Jobin-Yvon Horiba JY10000RF). Glancing-angle X-ray diffraction (GAXRD) analysis was applied to determine the crystallographic structure using Cu Ka radiation at an incidence angle of 2° (Panalytical X’Pert Pro). The coating structure was studied by scanning electron microscopy (SEM) on fracture cross-sections (Zeiss evo50). Micro-hardness measurements
were performed on polished samples with a computer-controlled micro-hardness tester (Fischerscope H100C) using a Vickers indenter at a load of 20 mN, at least 10 indentations and the BK7 tip shape correction [18]. 3. Results The influence of the AlCl3 partial pressure p(AlCl3) on the coating composition is shown in Fig. 1a and b for Tdep of 550 and 600 °C, respectively. From 550 to 600 °C, the concentration of reactive gases, i.e. AlCl3, TiCl4, and NH3 was reduced to improve the growth rate uniformity, but the ratio between the gases was kept constant. At 550 °C, the Al-content increases with p(AlCl3) from 0 at.% (i.e. TiN) up to 13 at.% Al. Associated with Al, the Cl-concentration rises from 3 to 7 at.%. Besides Cl, relatively high O-impurities have been observed. The Al-content increases at Tdep of 600 °C more intensely with p(AlCl3), but less steadily compared to 550 °C and the highest Al-content can be observed at the lowest p(AlCl3). Regarding the impurities, the Cl-concentration increases with p(AlCl3) up to 6 at.%, while the O-incorporation decreases. The temperature influence was examined between 600 and 700 °C using a constant p(AlCl3) of 62 Pa. As Tdep rises from 600 to 650 and 700 °C, the Al-content increases from 12 to 29 and finally 36 at.% Al. Simultaneously, the Cl- and O-concentration decreases from 5 to 3 at.% and 5 to 4 at.%, respectively. GAXRD patterns of six representative Ti–Al–N coatings within the investigated range (0–36 at.% Al) are exhibited in Fig. 2. The TiN coating (Tdep = 550 °C) shows a preferred (1 0 0) orientation (2H = 42.93°) and a considerably lower lattice parameter of a = 0.4210 nm compared to stoichiometric bulk material (a0 = 0.4240 nm). For 13 at.% Al (Tdep = 550 °C) and 19 at.% Al (Tdep = 600 °C), no additional phases can be detected and we refer to these coatings as fcc-Ti1–xAlxN. An Al-content of 20 at.%
Fig. 1. Coating composition in dependence on the AlCl3 partial pressure p(AlCl3) for deposition temperatures of (a) 550 °C and (b) 600 °C.
J. Wagner et al. / International Journal of Refractory Metals & Hard Materials 26 (2008) 563–568
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Fig. 2. GAXRD patterns of Ti–Al–N coatings with different Al-contents on cemented carbide substrates.
(Tdep = 600 °C) results in additional formation of fcc-AlN. Furthermore, at 24 at.% Al the first signs of hcp-AlN can be observed. High deposition temperatures and Al-contents, respectively, cause predominantly the formation of hcpAlN in favour of fcc-Ti1–xAlxN and fcc-AlN as can be seen for 36 at.% Al (Tdep = 700 °C) in Fig. 2. The (1 0 0) orientation of the fcc phase remains with increasing Al-incorporation, where the peaks show substantial broadening and the position approaches the standard value of TiN as the Alcontent increases. SEM micrographs of fracture cross-sections show that the micro-structure changes from columnar up to 13 at.% Al (Fig. 3a) to a dense, glassy morphology for 24 at.% Al (Fig. 3b) and to an agglomerate-like one for 36 at.% Al (Fig. 3c). The dependency of the micro-hardness HUpl on the Al-content is shown in Fig. 4. The coatings are distinguished according to the phases detected by GAXRD. HUpl decreases from 20 GPa for TiN to 13 GPa as the Al-concentration increases to 19 at.% Al. A higher content of 24 at.% consequences an enhanced micro-hardness of 20 GPa. Micro-hardness values could not be obtained for higher Al-concentrations since these coatings failed during the polishing process.
4. Discussion
Fig. 3. SEM fracture cross-sections of Ti–Al–N coatings containing: (a) 13 at.%; (b) 24 at.%; and (c) 36 at.% Al on cemented carbide substrates.
The AlCl3 partial pressure and, in particular, the deposition temperature have shown to remarkably influence the coating composition (Fig. 1), where both high temperatures and p(AlCl3)-values raise the Al-incorporation. The increasing Cl-concentration with p(AlCl3) indicates that Cl results mainly from an incomplete dissociation of the more stable AlCl3 and less from TiCl4. Higher temperatures enhance the reactivity of AlCl3 and consequence
higher Al- and lower Cl-contents in the coatings [1,14]. The relatively high O-contents detected may originate from residual oxygen in the deposition system, the high affinity of Cl-containing coatings to oxidation [19], and an O-enrichment at the surface (evidenced by GDOES depth-profiling) indicates oxidation at the end of the deposition process. A decreasing oxygen contamination with
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Fig. 4. Influence of Al-content on coating micro-hardness HUpl.
increasing Al-content can be noticed suggesting the formation of a protective alumina layer. GAXRD analysis indicates the formation of single-phase fcc-Ti1–xAlxN coatings up to 19 at.% Al (cf. Fig. 2). The coatings exhibit a preferred (1 0 0) orientation frequently found for TiN coatings using TiCl4 and NH3 [20,21]. Columnar micro-structures can be observed for low Al-contents, whereas higher concentrations result in grain refinement till glassy micro-structures are obtained. The decreasing grain size can be also detected in an enhanced XRD peak broadening. In contrary to PVD and plasmaassisted CVD fcc-Ti1–xAlxN coatings [1,22], the microhardness decreases with increasing Al-content up to 19 at.% (see Fig. 4). The simultaneously increasing Cl-content may lower the micro-hardness of these coatings [1,23]. Furthermore, low deposition temperatures are assumed to promote the formation of amorphous phases like a-AlN, which can not be excluded from XRD measurements [24]. Both Cl [23] and a-AlN could also contribute to the observed grain refinement due to segregation and formation of grain boundary phases, respectively, and thus hindered coalescence. The incorporation of Al into fcc-TiN can not be demonstrated from the dependency of the lattice parameter on the Al coating concentration. Generally, substitution of Ti by smaller Al atoms in the fcc-TiN structure gradually decreases the lattice parameter [1,15,22]. Here, the opposite trend can be observed and the lattice parameter increases with the Al-concentration. Several factors influence the lattice parameter and to elucidate an unambiguous correlation is difficult. The lattice parameter of TiN – ‘starting point’ for the following considerations – is already too low compared to the bulk value, as indicated by the peak shift shown in Fig. 2. From the low thermal expansion coefficient of cemented carbide, residual tensile stresses can be expected to arise in the coatings reducing the measured lattice parameter. The observed purple colour and the low lattice parameter of TiN point towards a non-stoichiometric compound TiNx (0.5 6 x 6 1.16), particularly an excess of N [25]. The vacancy structure of TiNx is stable over a wide range and causes lower lattice parameters for both
over- and under-stoichiometric coatings [26]. The measured excess of N relative to Ti (Fig. 1) indicates an over-stoichiometric TiNx (x > 1) phase, where the partially unoccupied Ti-sites cause shrinkage of the lattice. Low temperatures seem to favour the formation of non-stoichiometric TiNx; especially, under-stoichiometric TiNx has been found in many studies [20,21], whereas over-stoichiometric TiNx has been rarely mentioned [14,27]. The Cl- and O-impurities also affect the lattice parameter. Interstitial incorporation of large Cl atoms instead of N will cause an expansion of the lattice [28], while the substitution by O will shrink the lattice [29]. For the fcc-Ti1–xAlxN coatings grown in this work, it might be possible that Al takes the position of the vacant Ti-sites of the TiNx lattice, and as a result, the lattice increases with Al-incorporation. Between 20 and 24 at.% Al, further phases have been detected besides fcc-Ti1–xAlxN. The GAXRD patterns exhibit reflections of fcc-AlN at 20 at.% Al. The first indications of hcp-AlN can be observed at 24 at.% Al (see Fig. 2). In the following, these phases are denominated as fcc-AlN and hcp-AlN, respectively; although, Ti-incorporation can not be excluded. The additional phases cause fine-grained micro-structures and, compared to singlephase fcc-Ti1–xAlxN coatings, an enhanced micro-hardness (see Fig. 4). Metastable fcc-AlN has been detected during the decomposition of PVD fcc-Ti1–xAlxN coatings, where the supersaturated solution undergoes spinodal decomposition into coherent fcc-TiN and fcc-AlN domains in an fccTi1–xAlxN matrix. The system achieves equilibrium by separation of excess atoms into different phases. Finally, the metastable phases fcc-Ti1–xAlxN and fcc-AlN transform into stable fcc-TiN and hcp-AlN as the temperature increases. The indirect decomposition involving fcc-AlN reduces the energy barrier for hcp-AlN nucleation [4,30]. Here, a similar process could be responsible. Excessive Al-incorporation could destabilise the Ti-rich fccTi1–xAlxN and causes the precipitation of fcc-AlN during the deposition process. Surface-initiated spinodal decomposition of Ti1–xAlxN has already been observed to take place during growth of PVD fcc-Ti1–xAlxN, to form a rod-like nano-structure of fcc-TiN and fcc-AlN domains with a period of 2–3 nm [31]. As a result, the Ti-concentration in the fcc-Ti1–xAlxN matrix increases, which implies a reduction of the free energy [30]. Phase separation in the opposite way has been noticed in thermal CVD fccTi1–xAlxN coatings studies by Endler et al. [16]. There, an increase of Tdep resulted in additional formation of TiN besides high Al-containing fcc-Ti1–xAlxN (x = 0.8– 0.9). It can be expected that the high Al-content causes precipitation of TiN within the Al-supersaturated matrix, which reduces the excess Ti-concentration and, thus, the free energy of the fcc-Ti1–xAlxN phase. Al-contents above 24 at.% and high deposition temperatures, respectively, cause the formation of mainly hcpAlN at the expense of the fcc phases, since high temperatures favour the thermodynamically stable phases. Also, the XRD reflections of TiN approach the standard position
J. Wagner et al. / International Journal of Refractory Metals & Hard Materials 26 (2008) 563–568
with increasing temperature (Fig. 2). The formation of hcpAlN is disadvantageous for the coating properties, i.e. the micro-structure changes from dense, glassy to agglomerate-like (cf. Fig. 3b and c) and the micro-hardness decreases (cf. Fig. 4). 5. Conclusions Ti–Al–N coatings were deposited between 550 and 700 °C in an industrial thermal CVD system using TiCl4, AlCl3, and NH3. The Al-content increased with AlCl3concentration and deposition temperature. High temperatures enhance dissociation of metal chlorides, in particular of AlCl3 and raise the Al-content, while the incorporated Cl is reduced. Up to 19 at.% Al, only one face-centered cubic (fcc) phase has been detected. From the decreasing micro-hardness and the rising lattice parameter with Alincorporation, formation of a single-phase fcc-Ti1–xAlxN structure can not be unambiguously concluded, particularly, since possible amorphous phases like a-AlN are not gathered by X-ray diffraction. However, metastable fcc-AlN has been detected at 20 at.% and the first evidences of stable hexagonal close-packed (hcp) AlN at 24 at.% Al. The appearance of fcc-AlN is an indication for fcc-Ti1–xAlxN formation below 20 at.% Al and could result from spinodal decomposition of supersaturated fcc-Ti1–xAlxN. Al addition also changes the coating micro-structure from columnar to fine-grained or even glassy. While in general coatings show micro-hardness values decreasing with increasing Al (and Cl and O) content, the coating with 24 at.% Al yielded a relatively high micro-hardness of 20 GPa. High deposition temperatures and Al-concentrations up to 36 at.%, respectively, initiate predominantly the formation of hcp-AlN and TiN. These coatings yield an agglomerate-like morphology and unfavourable mechanical properties. This study reveals the feasibility to deposit Ti–Al–N coatings by thermal CVD in an industrial deposition unit and provides a basis for further developments with regard to coated cutting tools. Acknowledgments Financial support by the Austrian Kplus Competence Center Programme is gratefully acknowledged. The authors are grateful to K. Gigl for conducting the deposition processes and to G. Hawranek for SEM investigations. Special thanks are due to Dr. P.H. Mayrhofer for valuable discussions. References [1] Kim KH, Lee SJ. Comparative studies of TiN and Ti1 x Alx N by plasma-assisted chemical vapor deposition using a TiCl4/AlCl3/N2/ H2/Ar gas mixture. Thin Solid Films 1996;283:165–70. [2] Knotek O, Mu¨nz W-D, Leyendecker T. Industrial deposition of binary, ternary, and quaternary nitrides of titanium, zirconium, and aluminium. J Vac Sci Technol A 1987;5:2173–9.
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[3] Mu¨nz W-D. Titanium aluminum nitride films: a new alternative to TiN coatings. J Vac Sci Technol A 1986;4:2717–25. [4] Mayrhofer PH, Ho¨rling A, Karlsson L, Sjo¨le´n J, Larsson T, Mitterer C, et al. Self-organized nanostructures in the Ti–Al–N system. Appl Phys Lett 2003;83:2049–51. [5] Holleck H. Metastable coatings - prediction of composition and structure. Surf Coat Technol 1988;36:151–9. [6] Ho¨rling A, Hultman L, Ode´n M, Sjo¨le´n J, Karlsson L. Mechanical properties and machining performance of Ti1 x Alx N-coated cutting tools. Surf Coat Technol 2005;191:384–92. [7] Rebenne HE, Bhat DG. Review of CVD TiN coatings for wearresistant applications: deposition processes, properties and performance. Surf Coat Technol 1994;63:1–13. [8] Yoshikawa N. Review on processing, micro-structure and properties of vapour-deposited TiN and Ti-Al-N films. Recent Res Devel Mater Sci 2002;3:17–30. [9] Gilles S, Bourhila N, Ikeda S, Bernard C, Madar R. Deposition of (Ti,Al)N thin films by organometallic chemical vapor deposition: thermodynamic predictions and experimental results. Surf Coat Technol 1997;94–95:285–90. [10] Ikeda S, Gilles S, Chenevier B. Electron microscopy analysis of the microstructure of Ti1 x Alx N alloy thin films prepared using a chemical vapour deposition method. Thin Solid Films 1998;315:257–62. [11] Kurtz SR, Gordon RG. Chemical vapor deposition of titanium nitride at low temperatures. Thin Solid Films 1986;140:277–90. [12] Kim HJ, Egashira A, Komiyama H. Molecular size and its temperature dependence of growth species in chemical vapor deposition of aluminum nitride. J Chem Vap Dep 1992;1:20–41. [13] Yoshikawa N, Aoki M, Kikuchi A, Taniguchi S. Processing, microstructure and properties of chemical vapor deposited TiN/AlN films. Int J Mater Prod Technol 2001;16:109–15. [14] Anderbouhr S, Blanquet E, Ghetta V, Bernard C. TiN and (Ti1 x Alx)N films obtained by a low pressure chemical vapor deposition process. In: Allendorf MD, Bernard C, editors. Proc 14th Int conf chem vap dep and EUROCVD XI, Pennington: The Electrochemical Society; 1997. p. 356–63. [15] Anderbouhr S, Ghetta V, Blanquet E, Chabrol C, Schuster F, Bernard C, Madar R. LPCVD and PACVD (Ti,Al)N films: morphology and mechanical properties. Surf Coat Technol 1999;115:103–10. [16] Endler I, Herrmann M, Naupert M, Pitonak R, Ruppi S, Schneider M, van den Berg H, Westphal H. Aluminum-rich Ti1-xAlxN coatings by CVD. In: Proc euro PM2006. Shrewsbury: European Powder Metallurgy Association; 2006. p. 219–24. [17] Schintlmeister W, Wallgram W, Kanz J, Gigl K. Cutting tool materials coated by chemical vapour deposition. Wear 1984;100:153–69. [18] Heermant C, Dengl D. On the estimation of mechanical properties using the universal hardness test. Materialpru¨fung 1996;38:374–8. [19] Aizawa T, Akhadejdamrong T, Mitsuo A. Self-lubrication of nitride ceramic coating by the chlorine ion implantation. Surf Coat Technol 2004;177–178:573–81. [20] Huang HH, Hon MH. Effect of N2 addition on growth and properties of titanium nitride films obtained by atmospheric pressure chemical vapor deposition. Thin Solid Films 2002;416:54–61. [21] Ramanuja N, Levy RA, Dharmadhikari SN, Ramos E, Pearce CW, Menasian SC, Schamberger PC, Collins CC. Synthesis and characterization of low pressure chemically vapor deposited titanium nitride films using TiCl4 and NH3. Mater Lett 2002;57:261–9. [22] PalDey S, Deevi SC. Single layer and multilayer wear resistant coatings of (Ti,Al)N: a review. Mater Sci Eng A 2003;342:58–79. [23] Rie K-T, Gebauer A, Woehle J. Investigation of PA-CVD of TiN: relations between process parameters, spectroscopic measurements and layer properties. Surf Coat Technol 1993;60:385–8. [24] Liu YJ, Kim HJ, Egashira Y, Kimura H, Komiyama H. Using simultaneous deposition and rapid growth to produce nanostructured
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International Journal of Refractory Metals & Hard Materials 26 (2008) 563–568 www.elsevier.com/locate/IJRMHM
Deposition of Ti–Al–N coatings by thermal CVD J. Wagner a, V. Edlmayr a, M. Penoy b, C. Michotte b, C. Mitterer c,*, M. Kathrein d a
Materials Center Leoben Forschung GmbH, Roseggerstrasse 12, A-8700 Leoben, Austria b CERATIZIT Luxembourg S.a`.r.l., L-8201 Mamer, Luxembourg c Department of Physical Metallurgy and Materials Testing, University of Leoben, Franz-Josef-Strasse 18, A-8700 Leoben, Austria d CERATIZIT Austria GmbH, A-6600 Reutte, Austria Received 13 December 2007; accepted 17 January 2008
Abstract This study investigates the feasibility to deposit metastable Ti1–xAlxN coatings by thermal chemical vapour deposition (CVD) from the TiCl4–AlCl3–NH3–Ar system in an industrial CVD unit. The chemical composition and phase constitution was altered using different AlCl3 contents and deposition temperatures. According to X-ray diffraction analysis, single-phase face-centered cubic (fcc) structures have been obtained up to 19 at.% Al, where the incorporation of Al decreases grain size and micro-hardness. In addition, fcc-AlN has been detected between 20 and 24 at.% Al causing an enhanced micro-hardness and extremely fine structures. High temperatures and Al-contents produce coatings of poor mechanical properties consisting of TiN and hexagonal close-packed AlN. Ó 2008 Elsevier Ltd. All rights reserved. Keywords: Ti–Al–N; CVD; Structure; Hardness; Aluminium nitride
1. Introduction The deposition of Ti1–xAlxN (0 6 x 6 1) coatings by plasma-assisted chemical vapour deposition (CVD) [1] and physical vapour deposition (PVD) [2] techniques is well-established in cutting tool industry. High hardness and oxidation resistance enable the application in severe machining processes like high-speed and dry cutting [3]. The excellent properties result from the substitution of Ti by Al atoms in the face-centered cubic (fcc) TiN lattice, which improves both mechanical properties and thermal stability of TiN [4]. Although the mutual solubility of fcc-TiN and hexagonal close-packed (hcp) AlN is negligible in thermodynamic equilibrium, the non-equilibrium nature of plasma-assisted processes causes an extended solubility and the formation of metastable phases [5]. From the resulting supersaturated phases, fcc-Ti1–xAlxN (x < 0.6– 0.7) is ideal for cutting applications, while the presence of
*
Corresponding author. E-mail address: [email protected] (C. Mitterer).
0263-4368/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2008.01.003
hcp-AlxTi1–xN is widely accepted to deteriorate the performance [6]. The deposition of TiN for wear-resistant applications by thermal CVD is a conventional process since many years [7]. However, the incorporation of Al and formation of fcc-Ti1–xAlxN coatings without plasma assistance is challenging [8]. Single-phase fcc-Ti1–xAlxN (x < 0.57) coatings have been obtained by thermal CVD using metal-organic precursors [9,10]. Disadvantageous of metal-organic precursors are their high costs, and the low deposition temperatures may cause an insufficient adhesion. Most frequently used in cutting tool industry are the less expensive metal chlorides, e.g. TiCl4 together with N2 and H2 is a customary precursor system for the deposition of TiN. However, the incorporation of Al into TiN using AlCl3, N2, and H2 is thermodynamically not feasible. The utilization of NH3 as nitrogen source and reducing agent instead of N2 and H2 enables the formation of both TiN and AlN within a comparable temperature range [8]. Most studies consider the deposition of either single-phase coatings of fcc-TiN [11] and hcp-AlN [12], respectively, or fcc-TiN/hcp-AlN composites [13], but only few investigations report on
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fcc-Ti1–xAlxN deposits. Metastable fcc coatings have been deposited using TiCl3 and AlCl3 from in-situ chlorination of TiAl-alloys with Cl2 [14,15]. This process requires sophisticated equipment and bears the risk of preferred evaporation of a constituent. Recently, fcc-Ti1–xAlxN (x < 0.9) coatings have been deposited from TiCl4 and AlCl3 together with NH3 and H2 in a laboratory scale unit [16]. Here, we explore the feasibility to deposit Ti–Al–N coatings by thermal CVD, in particular fcc-Ti1–xAlxN, using the conventional precursors TiCl4 and AlCl3 together with NH3 in an industrial deposition plant. The Ti–Al–N coatings have been grown at different temperatures and AlCl3contents. The coatings have been characterised with respect to composition, structure, morphology, and hardness. 2. Experimental The deposition experiments were performed in an industrial vertical hot-wall CVD system (inner diameter of the reactor 205 mm, height 300 mm, usable for up to 1000 cemented carbide indexable inserts) at atmospheric pressure [17] using TiCl4, AlCl3, NH3, and Ar. AlCl3 was generated by reaction between HCl and aluminium chips. TiCl4 was provided by passing Ar through liquid TiCl4. The purity of the employed gases and reactants was better than 99.999%. The deposition unit was evacuated to remove residual gases and leak tested prior to all deposition processes. During the heating step, the deposition chamber was rinsed with H2 to remove organic contaminations and oxide layers on the substrates. Residual reactants and byproducts were removed at the end of the deposition process by rinsing the chamber with H2. The coating composition was adjusted by altering the deposition temperature and the AlCl3 partial pressure p(AlCl3). At a substrate temperature of 550 °C, p(AlCl3) was varied between 0 and 246 Pa, while other parameters were kept constant. The experiments between 600 and 700 °C were performed with half the concentration of reactive gases (i.e., AlCl3, TiCl4, and NH3) for improved growth rate homogeneity. Thus, p(AlCl3) was altered between 0 and 123 Pa at 600 °C. For the depositions at 650 and 700 °C, p(AlCl3) was set to 62 Pa. The coating deposition was performed on cemented carbide specimens (77 wt.% WC, 11 wt.% Co, 4 wt.% TiC, 8 wt.% (Ta,Nb)C). The sample position was the same throughout all deposition runs. The deposition time was 3 h for deposition temperatures Tdep of 550 and 600 °C and 2 h at Tdep of 650 and 700 °C. The mean chemical composition of the coatings was measured from glow discharge optical emission spectroscopy (GDOES) depth-profiles (Jobin-Yvon Horiba JY10000RF). Glancing-angle X-ray diffraction (GAXRD) analysis was applied to determine the crystallographic structure using Cu Ka radiation at an incidence angle of 2° (Panalytical X’Pert Pro). The coating structure was studied by scanning electron microscopy (SEM) on fracture cross-sections (Zeiss evo50). Micro-hardness measurements
were performed on polished samples with a computer-controlled micro-hardness tester (Fischerscope H100C) using a Vickers indenter at a load of 20 mN, at least 10 indentations and the BK7 tip shape correction [18]. 3. Results The influence of the AlCl3 partial pressure p(AlCl3) on the coating composition is shown in Fig. 1a and b for Tdep of 550 and 600 °C, respectively. From 550 to 600 °C, the concentration of reactive gases, i.e. AlCl3, TiCl4, and NH3 was reduced to improve the growth rate uniformity, but the ratio between the gases was kept constant. At 550 °C, the Al-content increases with p(AlCl3) from 0 at.% (i.e. TiN) up to 13 at.% Al. Associated with Al, the Cl-concentration rises from 3 to 7 at.%. Besides Cl, relatively high O-impurities have been observed. The Al-content increases at Tdep of 600 °C more intensely with p(AlCl3), but less steadily compared to 550 °C and the highest Al-content can be observed at the lowest p(AlCl3). Regarding the impurities, the Cl-concentration increases with p(AlCl3) up to 6 at.%, while the O-incorporation decreases. The temperature influence was examined between 600 and 700 °C using a constant p(AlCl3) of 62 Pa. As Tdep rises from 600 to 650 and 700 °C, the Al-content increases from 12 to 29 and finally 36 at.% Al. Simultaneously, the Cl- and O-concentration decreases from 5 to 3 at.% and 5 to 4 at.%, respectively. GAXRD patterns of six representative Ti–Al–N coatings within the investigated range (0–36 at.% Al) are exhibited in Fig. 2. The TiN coating (Tdep = 550 °C) shows a preferred (1 0 0) orientation (2H = 42.93°) and a considerably lower lattice parameter of a = 0.4210 nm compared to stoichiometric bulk material (a0 = 0.4240 nm). For 13 at.% Al (Tdep = 550 °C) and 19 at.% Al (Tdep = 600 °C), no additional phases can be detected and we refer to these coatings as fcc-Ti1–xAlxN. An Al-content of 20 at.%
Fig. 1. Coating composition in dependence on the AlCl3 partial pressure p(AlCl3) for deposition temperatures of (a) 550 °C and (b) 600 °C.
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Fig. 2. GAXRD patterns of Ti–Al–N coatings with different Al-contents on cemented carbide substrates.
(Tdep = 600 °C) results in additional formation of fcc-AlN. Furthermore, at 24 at.% Al the first signs of hcp-AlN can be observed. High deposition temperatures and Al-contents, respectively, cause predominantly the formation of hcpAlN in favour of fcc-Ti1–xAlxN and fcc-AlN as can be seen for 36 at.% Al (Tdep = 700 °C) in Fig. 2. The (1 0 0) orientation of the fcc phase remains with increasing Al-incorporation, where the peaks show substantial broadening and the position approaches the standard value of TiN as the Alcontent increases. SEM micrographs of fracture cross-sections show that the micro-structure changes from columnar up to 13 at.% Al (Fig. 3a) to a dense, glassy morphology for 24 at.% Al (Fig. 3b) and to an agglomerate-like one for 36 at.% Al (Fig. 3c). The dependency of the micro-hardness HUpl on the Al-content is shown in Fig. 4. The coatings are distinguished according to the phases detected by GAXRD. HUpl decreases from 20 GPa for TiN to 13 GPa as the Al-concentration increases to 19 at.% Al. A higher content of 24 at.% consequences an enhanced micro-hardness of 20 GPa. Micro-hardness values could not be obtained for higher Al-concentrations since these coatings failed during the polishing process.
4. Discussion
Fig. 3. SEM fracture cross-sections of Ti–Al–N coatings containing: (a) 13 at.%; (b) 24 at.%; and (c) 36 at.% Al on cemented carbide substrates.
The AlCl3 partial pressure and, in particular, the deposition temperature have shown to remarkably influence the coating composition (Fig. 1), where both high temperatures and p(AlCl3)-values raise the Al-incorporation. The increasing Cl-concentration with p(AlCl3) indicates that Cl results mainly from an incomplete dissociation of the more stable AlCl3 and less from TiCl4. Higher temperatures enhance the reactivity of AlCl3 and consequence
higher Al- and lower Cl-contents in the coatings [1,14]. The relatively high O-contents detected may originate from residual oxygen in the deposition system, the high affinity of Cl-containing coatings to oxidation [19], and an O-enrichment at the surface (evidenced by GDOES depth-profiling) indicates oxidation at the end of the deposition process. A decreasing oxygen contamination with
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Fig. 4. Influence of Al-content on coating micro-hardness HUpl.
increasing Al-content can be noticed suggesting the formation of a protective alumina layer. GAXRD analysis indicates the formation of single-phase fcc-Ti1–xAlxN coatings up to 19 at.% Al (cf. Fig. 2). The coatings exhibit a preferred (1 0 0) orientation frequently found for TiN coatings using TiCl4 and NH3 [20,21]. Columnar micro-structures can be observed for low Al-contents, whereas higher concentrations result in grain refinement till glassy micro-structures are obtained. The decreasing grain size can be also detected in an enhanced XRD peak broadening. In contrary to PVD and plasmaassisted CVD fcc-Ti1–xAlxN coatings [1,22], the microhardness decreases with increasing Al-content up to 19 at.% (see Fig. 4). The simultaneously increasing Cl-content may lower the micro-hardness of these coatings [1,23]. Furthermore, low deposition temperatures are assumed to promote the formation of amorphous phases like a-AlN, which can not be excluded from XRD measurements [24]. Both Cl [23] and a-AlN could also contribute to the observed grain refinement due to segregation and formation of grain boundary phases, respectively, and thus hindered coalescence. The incorporation of Al into fcc-TiN can not be demonstrated from the dependency of the lattice parameter on the Al coating concentration. Generally, substitution of Ti by smaller Al atoms in the fcc-TiN structure gradually decreases the lattice parameter [1,15,22]. Here, the opposite trend can be observed and the lattice parameter increases with the Al-concentration. Several factors influence the lattice parameter and to elucidate an unambiguous correlation is difficult. The lattice parameter of TiN – ‘starting point’ for the following considerations – is already too low compared to the bulk value, as indicated by the peak shift shown in Fig. 2. From the low thermal expansion coefficient of cemented carbide, residual tensile stresses can be expected to arise in the coatings reducing the measured lattice parameter. The observed purple colour and the low lattice parameter of TiN point towards a non-stoichiometric compound TiNx (0.5 6 x 6 1.16), particularly an excess of N [25]. The vacancy structure of TiNx is stable over a wide range and causes lower lattice parameters for both
over- and under-stoichiometric coatings [26]. The measured excess of N relative to Ti (Fig. 1) indicates an over-stoichiometric TiNx (x > 1) phase, where the partially unoccupied Ti-sites cause shrinkage of the lattice. Low temperatures seem to favour the formation of non-stoichiometric TiNx; especially, under-stoichiometric TiNx has been found in many studies [20,21], whereas over-stoichiometric TiNx has been rarely mentioned [14,27]. The Cl- and O-impurities also affect the lattice parameter. Interstitial incorporation of large Cl atoms instead of N will cause an expansion of the lattice [28], while the substitution by O will shrink the lattice [29]. For the fcc-Ti1–xAlxN coatings grown in this work, it might be possible that Al takes the position of the vacant Ti-sites of the TiNx lattice, and as a result, the lattice increases with Al-incorporation. Between 20 and 24 at.% Al, further phases have been detected besides fcc-Ti1–xAlxN. The GAXRD patterns exhibit reflections of fcc-AlN at 20 at.% Al. The first indications of hcp-AlN can be observed at 24 at.% Al (see Fig. 2). In the following, these phases are denominated as fcc-AlN and hcp-AlN, respectively; although, Ti-incorporation can not be excluded. The additional phases cause fine-grained micro-structures and, compared to singlephase fcc-Ti1–xAlxN coatings, an enhanced micro-hardness (see Fig. 4). Metastable fcc-AlN has been detected during the decomposition of PVD fcc-Ti1–xAlxN coatings, where the supersaturated solution undergoes spinodal decomposition into coherent fcc-TiN and fcc-AlN domains in an fccTi1–xAlxN matrix. The system achieves equilibrium by separation of excess atoms into different phases. Finally, the metastable phases fcc-Ti1–xAlxN and fcc-AlN transform into stable fcc-TiN and hcp-AlN as the temperature increases. The indirect decomposition involving fcc-AlN reduces the energy barrier for hcp-AlN nucleation [4,30]. Here, a similar process could be responsible. Excessive Al-incorporation could destabilise the Ti-rich fccTi1–xAlxN and causes the precipitation of fcc-AlN during the deposition process. Surface-initiated spinodal decomposition of Ti1–xAlxN has already been observed to take place during growth of PVD fcc-Ti1–xAlxN, to form a rod-like nano-structure of fcc-TiN and fcc-AlN domains with a period of 2–3 nm [31]. As a result, the Ti-concentration in the fcc-Ti1–xAlxN matrix increases, which implies a reduction of the free energy [30]. Phase separation in the opposite way has been noticed in thermal CVD fccTi1–xAlxN coatings studies by Endler et al. [16]. There, an increase of Tdep resulted in additional formation of TiN besides high Al-containing fcc-Ti1–xAlxN (x = 0.8– 0.9). It can be expected that the high Al-content causes precipitation of TiN within the Al-supersaturated matrix, which reduces the excess Ti-concentration and, thus, the free energy of the fcc-Ti1–xAlxN phase. Al-contents above 24 at.% and high deposition temperatures, respectively, cause the formation of mainly hcpAlN at the expense of the fcc phases, since high temperatures favour the thermodynamically stable phases. Also, the XRD reflections of TiN approach the standard position
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with increasing temperature (Fig. 2). The formation of hcpAlN is disadvantageous for the coating properties, i.e. the micro-structure changes from dense, glassy to agglomerate-like (cf. Fig. 3b and c) and the micro-hardness decreases (cf. Fig. 4). 5. Conclusions Ti–Al–N coatings were deposited between 550 and 700 °C in an industrial thermal CVD system using TiCl4, AlCl3, and NH3. The Al-content increased with AlCl3concentration and deposition temperature. High temperatures enhance dissociation of metal chlorides, in particular of AlCl3 and raise the Al-content, while the incorporated Cl is reduced. Up to 19 at.% Al, only one face-centered cubic (fcc) phase has been detected. From the decreasing micro-hardness and the rising lattice parameter with Alincorporation, formation of a single-phase fcc-Ti1–xAlxN structure can not be unambiguously concluded, particularly, since possible amorphous phases like a-AlN are not gathered by X-ray diffraction. However, metastable fcc-AlN has been detected at 20 at.% and the first evidences of stable hexagonal close-packed (hcp) AlN at 24 at.% Al. The appearance of fcc-AlN is an indication for fcc-Ti1–xAlxN formation below 20 at.% Al and could result from spinodal decomposition of supersaturated fcc-Ti1–xAlxN. Al addition also changes the coating micro-structure from columnar to fine-grained or even glassy. While in general coatings show micro-hardness values decreasing with increasing Al (and Cl and O) content, the coating with 24 at.% Al yielded a relatively high micro-hardness of 20 GPa. High deposition temperatures and Al-concentrations up to 36 at.%, respectively, initiate predominantly the formation of hcp-AlN and TiN. These coatings yield an agglomerate-like morphology and unfavourable mechanical properties. This study reveals the feasibility to deposit Ti–Al–N coatings by thermal CVD in an industrial deposition unit and provides a basis for further developments with regard to coated cutting tools. Acknowledgments Financial support by the Austrian Kplus Competence Center Programme is gratefully acknowledged. The authors are grateful to K. Gigl for conducting the deposition processes and to G. Hawranek for SEM investigations. Special thanks are due to Dr. P.H. Mayrhofer for valuable discussions. References [1] Kim KH, Lee SJ. Comparative studies of TiN and Ti1 x Alx N by plasma-assisted chemical vapor deposition using a TiCl4/AlCl3/N2/ H2/Ar gas mixture. Thin Solid Films 1996;283:165–70. [2] Knotek O, Mu¨nz W-D, Leyendecker T. Industrial deposition of binary, ternary, and quaternary nitrides of titanium, zirconium, and aluminium. J Vac Sci Technol A 1987;5:2173–9.
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