Development of nanostructures in thin films of transition metal nitrides and their properties

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1 Development of nanostructures in thin films of transition metal nitrides and their properties David Rafaja*, Christina Wüstefeld, Steffen Wolf, Milan Dopita, Dietrich Heger: Institute of Materials Science, TU Bergakademie Freiberg, Gustav-Zeuner-Str. 5, D Freiberg, Germany Michal Šíma: SHM Ltd., Prmyslová 3, CZ Šumperk, Czech Republic * Corresponding author. Tel.: , Fax: , rafaja@ww.tu-freiberg.de Abstract Hardness, phase composition and stress-free lattice parameters of the face centred cubic phases in the (M, Al) N thin film nanocomposites with M = Cr, Ti or Zr were investigated at different aluminium contents in the samples. The hardness was measured using the Oliver & Pharr method, the chemical composition using the electron probe microanalysis with wavelength-dispersive analysis of characteristic X-ray spectra. The phase composition of the samples and the stress-free lattice parameters of the fcc phases were obtained from the X-ray diffraction measurements performed in the glancing-angle diffraction geometry. In all three systems, the maximum hardness was observed in the samples containing two nanocrystalline phases, i.e. fcc-m 1-x Al x N and w-aln with the wurtzite crystal structure. For the fcc-m 1-x Al x N phases, the stability ranges and the dependences of the lattice parameters on the aluminium contents within these stability ranges were determined. The maximum extents of the lattice parameters within the fcc phases are discussed. Keywords: Ultra-hard thin film nanocomposites, cathodic arc evaporation, transition metal aluminium nitrides, phase stability, lattice parameters, hardness In the concept of thin film nanocomposites based on the work published in [1, 2], the hardness is improved mainly by the small crystallite size as proposed by Hall [3] and Petch [4], as a small crystallite size reduces typically the movement of the structure defects through the crystallites. For a further increase of the hardness of nanocrystalline materials, a reduction of the grain boundary sliding is needed that reduces the propagation of structure defects. The grain boundary sliding can be reduced by cohesion of neighbouring nanocrystallites [5]; this implies that the morphology of the internal interfaces plays a crucial role in the development of the hardness of nanocomposites [6]. The Hall-Petch like behaviour of the hardness was observed in numerous transition metal (mainly Ti or Cr) nitride thin film nanocomposites containing additionally aluminium and silicon [1, 2, 6 11]. In both, Ti-Al-Si-N and Cr-Al-Si-N systems the hardness of the thin film nanocomposites was found to be improved by coherency strains at the boundaries of crystallites having different crystal structures, i.e. fcc-m 1-x Al x N (M = Ti, Cr) and w-aln [8 11]. Thus, it can be expected that the phase composition of the nanocomposites will strongly affect their hardness. This hypothesis was verified experimentally in [12], where it was shown that the

2 maximum hardness in the (Ti, Al) N thin films lies near the aluminium concentration, at which the material decomposes into the nanocrystalline fcc-ti 1-x Al x N and w-aln domains. This contribution illustrates the correlation between the hardness of the Cr 1-x Al x N, Ti 1-x Al x N and Zr 1-x Al x N thin film nanocomposites and their phase composition. The thin film nanocomposites were deposited using the cathodic arc evaporation (CAE) in nitrogen atmosphere at the working pressure of 1.3 Pa from two separated cathodes [13]. The current on the Cr, Ti or Zr cathode was 80 A, the current on the Al cathode 120 A. The hardness and the phase composition of the thin films were related to their overall chemical composition that varied with the position of the samples in the deposition apparatus and that was measured using the electron probe microanalysis with wavelength dispersive analysis of characteristic X-ray spectra (EPMA/WDAXS). The phase composition was analysed with the aid of the glancing angle X-ray diffraction (GAXRD) performed on a D8 diffractometer (Bruker AXS). The diffractometer was equipped by a sealed X-ray tube with copper radiation, by a Goebel mirror in the primary beam, and a by Soller collimator and a LiF monochromator both located in the diffracted beam. Furthermore, the GAXRD patterns were used for determination of the stress-free lattice parameters by using the modified sin²ψ method [9]. For TiN, the Poisson ratio of 0.3 [14] was used. CrN and ZrN have shown a strong anisotropy of the lattice deformation, thus the lattice deformation was recalculated into the hard direction 100 as shown in [9]. Accordingly, the Poisson ratios for the crystallographic direction 100 were used for CrN and ZrN, i.e and 0.186, respectively [15]. The dependence of the stress-free lattice parameter on the overall chemical composition was employed as a complementary method for precise determination of the phase composition. Figure 1: Dependence of the hardness of the CAE thin films on the [M]/([M]+[Al]) ratio measured in the Cr 1-x Al x N, Ti 1-x Al x N and Zr 1-x Al x N thin films. The results of the hardness measurement performed according to the Oliver & Pharr method [16] are shown in Fig. 1. In all systems under study, the maximum hardness was found in a certain composition range that shifted to lower aluminium concentrations with increasing intrinsic lattice parameters in the respective binary nitride, a 0 (CrN) = nm, a 0 (TiN) = nm and a 0 (ZrN) = nm [17]. The dependence of the stress-free lattice parameters on the transition metal contents in the M 1-x Al x N nanocrystalline thin films is shown in Fig. 2. In all cases, the stress-free lattice parameter decreased with increasing aluminium contents. At low 2

3 aluminium contents, Vegard-like linear dependences of the stress-free lattice parameters on the transition metal contents were observed that can be expected if aluminium replaces the transition metal atoms in the cubic elementary cell. In [8], [9] and [11] it was shown with the aid of the combination of the phase analysis and the measurement of the stress-free lattice parameter that the linear dependence of a versus [M]/([M]+[Al]) is observed in samples containing a single phase. In samples containing two phases, i.e. fcc-(m, Al) N and w-aln, the aluminium contents in fcc-(m, Al) N is lower than the overall aluminium contents measured using EPMA/WDAXS, thus the stress-free lattice parameter is higher than the lattice parameter expected for the measured overall chemical composition. The linear parts of the dependences of the stress-free lattice parameters on the transition metal contents were described using the functions: a(cr 1-x Al x N) = [ (2) (1) x] nm, a(ti 1-x Al x N) = [ (2) (2) x] nm and a(zr 1-x Al x N) = [0.458(1) 0.026(1) x] nm. The parameters of these linear functions show that the decrease of the lattice parameter with increasing aluminium contents becomes faster for transition metal nitrides with larger intrinsic lattice parameters. Figure 2: Dependence of the stress-free lattice parameters on the [M]/([M]+ [Al]) ratio; M = Cr, Ti, Zr. Grey boxes indicate the chemical composition, for which the maximum hardness was observed (compare with Fig. 1). Solid lines show the Vegard-like dependences of the lattice parameters calculated as obtained for the samples containing a single phase. The comparison of the composition ranges, in which the maximum hardness was detected, with the phase composition indicated that the maximum hardness appears in thin film nanocomposites containing both fcc-(m, Al) N and w-aln. The composition ranges, in which the maximum hardness was observed, match with the composition ranges, in which the departure of the lattice parameters of the fcc-(m, Al) N from the Vegard-like dependence was found (Fig. 2). At higher Al contents, where w-aln becomes the dominant phase, the hardness of the thin films decreases (Fig. 1). The Vegard-like dependences of the lattice parameters on the Al contents are summarised in Fig. 3 for Cr 1-x Al x N, Ti 1-x Al x N and Zr 1-x Al x N together with the homogeneity ranges of these phases that were obtained from the combination of the XRD phase analysis and the precise measurement of the stress-free lattice parameters. The extent of the lattice parameters within the fcc-(m, Al) N phases remains almost the same in all systems under study, a/a This could be explained by a certain maximum intrinsic lattice strain that is due to the replacement of the transition metal atoms by Al. A consequence of the nearly constant 3

4 extent of the lattice parameters is that the stability range of the respective fcc phase becomes narrower with increasing slope of the dependence of the lattice parameter on the Al contents. Figure 3: Summary of the Vegard-like dependences of the lattice parameters for the fcc phases in the Cr 1-x Al x N, Ti 1-x Al x N and Zr 1-x Al x N thin films. The grey boxes point out the stability range of the respective fcc phase. Their horizontal size indicates the width of the stability range, their vertical size the maximum extent of the stress-free lattice parameters as observed in the (M, Al) N thin film nanocomposites deposited using CAE. As based on the above results, we can conclude that the phase composition is a very important microstructure parameter, which influences the hardness of the (Ti, Al) N, (Cr, Al) N and (Zr, Al) N thin film nanocomposites. For all three systems, the maximum hardness was observed in samples containing two crystalline phases in similar quantities. The hardness decreased towards the single fcc-phase samples as well as towards the samples with prevailing w-aln. Although Ti 1-x Al x N, Cr 1-x Al x N and Zr 1-x Al x N have different stability ranges, the extent of the respective lattice parameter is very similar in all three systems. Acknowledgements: The authors would like to thank the German Scientific Council (DFG) for supporting the project # RA 1050/9-1 and the Dr. Erich Krüger foundation of the TU Bergakademie Freiberg for supporting the project of the Freiberg High-Pressure Centre. References [1] S. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64. [2] S. Veprek, S. Reiprich, Li Shizhib, Appl. Phys. Lett. 66 (1995) [3] E.O. Hall, Proc. Phys. Soc., London, Sect. B 64 (1951) 747. [4] N.J. Petch, J. Iron Steel Inst. 174 (1953) 25. [5] S. Veprek, A.S. Argon, J. Vac. Sci. Technol. B 20 (2002) 650. [6] S. Veprek, M.G.J. Veprek-Heijman, Surf. Coat. Technol. 201 (2007) [7] P.H. Mayrhofer, G. Tischler, C. Mitterer, Surf. Coat. Technol (2001) 78. [8] D. Rafaja, A. Poklad, V. Klemm, G. Schreiber, D. Heger, M. Šíma, M. Dopita, Thin Solid Films 514 (2006) 240. [9] D. Rafaja, M. Dopita, M. Ržika, V. Klemm, D. Heger, G. Schreiber, M. Šíma, Surf. Coat. Technol. 201 (2006)

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