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X-ray absorption near-edge structure of hexagonal ternary phases in sputter-deposited TiAlN films
Journal of Alloys and Compounds 561 (2013) 87–94
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
X-ray absorption near-edge structure of hexagonal ternary phases in sputter-deposited TiAlN films R. Gago a,⇑, F. Soldera b, R. Hübner c, J. Lehmann c, F. Munnik c, L. Vázquez a, A. Redondo-Cubero d, J.L. Endrino a,e a
Instituto de Ciencia de Materiales de Madrid, Consejo Superior de Investigaciones Científicas, E-28049 Madrid, Spain Department of Materials Science & Engineering, Saarland University, D-66123 Saarbruecken, Germany Institute of Ion Beam Physics and Materials Research, Helmholtz-Zentrum Dresden-Rossendorf, D-01314 Dresden, Germany d Instituto Tecnológico e Nuclear, Instituto Superior Técnico, Universidade Técnica de Lisboa, 2686-953 Sacavém, Portugal e Abengoa Research S.L., c/Energía Solar 1, Palmas Altas, E-41014 Seville, Spain b c
a r t i c l e
i n f o
Article history: Received 12 December 2012 Received in revised form 21 January 2013 Accepted 23 January 2013 Available online 13 February 2013 Keywords: Nitride materials Vapour deposition Atomic scale structure NEXAFS/XANES
a b s t r a c t Titanium aluminium nitride (TiAlN) coatings have been grown by reactive (Ar/N2) direct-current magnetron sputtering from a Ti50Al50 compound target. The film composition has been quantified by ion beam analysis showing the formation of Al-rich nitrides (Ti/Al 0.3), with stoichiometric films for N2 contents in the gas mixture equal or above 25%. The surface morphology of the films has been imaged by atomic force microscopy, showing very smooth surfaces with roughness values below 2 nm. X-ray and electron diffraction patterns reveal that the films are nanocrystalline with a wurzite (w) structure of lattice parameters larger (2.5%) than those for w-AlN. The lattice expansion correlates with the Ti/Al ratio in stoichiometric films, which suggests the incorporation of Ti into w-AlN. The atomic environments around Ti, Al and N sites have been extracted from the X-ray absorption near-edge structure (XANES) by recording the Ti2p, Al1s and N1s edges, respectively. The analysis of the XANES spectral lineshape and comparison with reported theoretical calculations confirm the formation of a ternary hexagonal phase. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Transition metal nitride materials have been commercially available for many decades as hard, wear-resistant, and decorative coatings, as well as electrical contacts and conductive diffusion barriers in electronic devices [1]. Among them, titanium nitride (TiN) is widely used as protective coating material in cutting tools, mechanical components and medical implants due to its high hardness, non-toxicity, low friction, and chemical inertness [2]. Ternary compounds within the Ti–Al–N system have been extensively studied as an alternative to TiN overcoats for high-temperature applications since Al improves the wear and oxidation resistance, as well as the thermal stability of the coated system [3]. In the case of ternary alloys, a crucial aspect to tune the final properties is an accurate control of both the composition and structure of the thin films. Hence, the production of ternary TiAlN compounds has been mainly focused on increasing the Al content in the cubic TiN (c-TiN) structure (solid solution) to maximize oxidation resistance while preventing the segregation of binary phases [4]. However, c-TiN can only accommodate up to a certain amount of AlN molar fractions (x) ⇑ Corresponding author at: Instituto de Ciencia de Materiales de Madrid, Consejo Superior de Investigaciones Científicas, E-28049 Madrid, Spain. Tel.: +34913349090. E-mail address: [email protected] (R. Gago). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.01.130
within a metastable N-saturated solid solution (c-Ti1xAlxN) [5]. As the solubility limit is exceeded, the film structure undergoes a spinodal decomposition (i.e., c-Ti1xAlxN domains with different composition) [6] or phase segregation of c-Ti1xAlxN and (hexagonal) wurzite aluminium nitride (w-AlN) [7]. Theoretically, the critical AlN molar fraction is predicted to be x 0.7 [8]. Experimentally, a broad range of solubility limits have been published (0.4 < x < 0.9) [9], which could be understood based on the multiple growth conditions employed. Contrary to c-Ti1xAlxN, the formation of hexagonal ternary single-phases of w-Ti1xAlxN (that is, the incorporation of Ti into w-AlN) has not been studied in detail so far. Such phases are softer than the cubic counterpart but present an excellent oxidation resistance [10]. Hence, they can be suitable candidates for certain technological applications. Further, the study of w-Ti1xAlxN phases could be relevant to understand the formation and electronic structure of nanolaminated ternary materials in the Ti–Al– N system with hexagonal structure [11]. In this case, it has been predicted [12] that the wurzite structure can only accommodate Ti contents up to 20 at.% (x 0.6), being unstable above this threshold. X-ray diffraction (XRD) is the most common technique used to study the microstructure of Ti–Al–N coatings. However, XRD probes only the crystallized domains and, hence, information from the eventual participation of amorphous nitrides or grain boundaries
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is lacking [13]. Further, the position of the Bragg reflections can overlap for different structures or be affected by other issues such as lattice strain. Complementary, the near-edge structure (NES) extracted from electron energy loss (ELNES) or X-ray absorption (XANES) measurements can be used to solve complex systems and discern the coexistence of different phases, even in amorphous or disordered structures. Typically, XANES presents a better spectral resolution than ELNES (0.1 eV instead of 1 eV) although it requires monochromatic X-rays from a synchrotron facility [14]. These techniques probe the density of unoccupied states and yield local order information with element selectivity by recording specific edges. In this way, the simultaneous analysis of the characteristic absorption edge(s) for each element can be unequivocally used for phase identification (distinct spectral features are used as phase fingerprints). Despite the potential of this analysis, there are only few XANES reports addressing the formation [15–17] and thermal stability [18] of ternary TiAlN compounds. It should be noted that c-AlN is a metastable phase and it is only produced under highpressure [19], whereas TiN is mechanically unstable in the fourcoordinated wurzite structure [12]. Therefore, identification of cubic or hexagonal ternary environments by XANES may be complex due to the lack of reference nitrides with Al and Ti sites occupied in the rock-salt or hexagonal structure, respectively. Nevertheless, XANES spectra of c-AlN have been successfully recorded by stabilizing this phase in short-period TiN/AlN multilayers under compressive stress [20] and upon segregation in low carbon aluminiumkilled steels [21]. Alternatively, theoretical modelling of ELNES or XANES spectra can be used to predict and understand the bonding behaviour of ternary TiAlN phases with either hexagonal or cubic structure [12]. However, experimental evidence to support such predictions is still needed, especially, in the case of hexagonal phases. In this work, the formation of hexagonal TiAlN ternary phases is identified by the combination of XRD and XANES. Soft X-rays are used to study the atomic arrangements around Ti, Al and N by recording the Ti2p, Al1s and N1s XANES spectra, respectively. The presence of distinct spectral features in the XANES measurements with respect to binary nitrides and comparison with reported theoretical results support the formation of a ternary phase. To the best of our knowledge, XANES of w-Ti1xAlxN phases has not been reported experimentally so far. This information will be relevant to identify the formation of ternary phases in future work and will provide experimental data for further modelling and understanding the electronic structure of ternary nitride systems. 2. Materials and methods 2.1. Sample preparation TiAlN films were grown on Si(1 0 0) substrates by reactive direct-current (DC) magnetron sputtering from a 3’’ Ti50Al50 (in at.%) composite target. The cathode was facing the grounded substrates at a working distance of 15 cm. The base and working pressure were 104 and 0.2 Pa, respectively. Different Ar (99.9995% purity grade) and N2 (99.9995% purity grade) contents in the gas mixture were used to influence the film properties. For this purpose, the individual Ar (UAr) and N2 (UN2) flows were adjusted while keeping the total flow (pressure) constant at 20 standard cubic centimeter per minute (sccm). In this way, the growth has been parameterized as a function of the N2 content in the gas mixture, [N2] = UN2/ (UAr + UN2). For plasma generation, a DC voltage with an overall power of 200 W was applied to the cathode. The depositions were carried out for 1 h at room temperature. Additionally, a thick sample (>1 lm) was grown for 6 h with [N2] = 50% to study the microstructure evolution with thickness. In this case, the film was grown onto a cemented carbide insert to improve the adhesion with respect to Si(1 0 0) substrates. 2.2. Sample characterisation The thickness of the TiAlN films was determined with a Dektak-150 (VeecoÒ) mechanical profiler by measuring the step height created by partially masking the substrate during growth. The composition of the layers was determined by
heavy-ion elastic recoil detection analysis (ERDA) with a 35 MeV Cl7+ beam impinging at 15° with respect to the sample surface. The scattered ions and recoils were detected with a Bragg ionisation chamber (BIC) located at a scattering angle of 31° that discriminates the detected particle according to their atomic number (Z). The measured energy spectrum of each element was converted into compositional depth profiles by means of the analysis program WiNDF [22] using the stopping power data from Ziegler et al. [23]. The resulting surface morphology was imaged by atomic force microscopy (AFM) operating in the dynamic mode with an Agilent PicoPlus 5500. AFM probes with a nominal radius of 8 nm and opening angle smaller than 52° were used. Image editing and analysis was done with the Gwyddion freeware package [24]. The microstructure of the samples was examined by glancing-angle XRD measurements using a D5000 (BRUKER AXS) diffractometer with Cu Ka radiation (k = 1.5418 Å). The data were collected with an incidence angle of 1° within the scattering range, 2h, of 30–80°. Complementary, the bonding structure with element sensitivity to Ti, Al and N sites was studied by XANES with soft X-rays at the Canadian Light Source (CLS) synchrotron using the SGM beamline, which provides an energy resolution of 0.1 eV. The measurements were performed in the total electron yield (TEY) mode with the sample normal positioned parallel to the X-ray beam. Finally, the microstructural evolution was examined by crosssectional transmission electron microscopy (TEM) with an image-corrected TITAN 80-300 (FEI) microscope operated at 300 kV. For this purpose, a slice from the thick TiAlN film was prepared in a focused ion beam and scanning electron microscope dual beam system (FEI Helios 600) by using the ‘‘in situ’’ lift-out technique. The foil was thinned to electron transparency, first, with an acceleration voltage of 30 kV and, later, with 5 kV (for at least 2 min in each side) to minimize any possible preparation artifacts.
3. Results and discussion 3.1. Thickness, composition and surface morphology The thickness of the TiAlN films as a function of N2 content in the gas mixture, [N2], is shown in Table 1. Clearly, the thickness decreases progressively with [N2] or, equivalently, the deposition rate drops from 7 to 2 nm/min. This trend is the result of target poisoning (nitride formation) and the transition toward the sputtering compound mode by increasing the reactive gas content [25]. As an example, the ERDA-BIC data from a TiAlN film grown with [N2] = 25% is shown in Fig. 1a. The signals from the elements in the film and the scattered Cl projectiles have been labelled in the graph. The potential of this analysis is that well-resolved signals are obtained, specially, for light elements. In addition, heavier elements such as Ti provide double information from Ti recoils and Cl projectiles scattered from Ti target atoms. For each detected particle, lower energies imply scattering events deeper in the sample. The energy onset for elements at the surface is given by the kinematic factor of the collision (this onset is higher for heavier elements). Note that oxygen is only detected at the surface as an indication of surface contamination and/or oxidation, and at the substrate/film interface from the native oxide of the Si(1 0 0) substrate. In addition, a negligible concentration of C is detected, showing the high purity grade of the process. In the ERDA-BIC map, the signal from each element can be projected independently to the energy (abscissa) axis and, in this way, the element depth profile can be derived. As an example, the compositional profile derived from the ERDA-BIC data of Fig. 1a is plotted in panel b. It can be seen that a stoichiometric nitride is obtained with a composition close to Ti0.25Al0.75N. The average film composition has been quantified by integrating the compositional profiles derived from ERDA-BIC over the whole film thickness. The resulting values are included in Table 1 showing an increase of the N content with [N2]. A sub-stoichiometric nitride is obtained for the lowest value of [N2] = 10%, which means that part of the metal atoms do not react to form a nitride due to deficient N supply. On the contrary, stoichiometric nitrides (N = 50 at.%) are formed for [N2] P 25%, with slightly over-stoichiometric values for [N2] P 50%. In this case, the slight excess of N is probably incorporated in the form of embedded N2 molecules, as observed for the growth of TiN films [26]. Note that ERDA-BIC is
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Table 1 Film thickness (Th), composition, surface roughness (r), wurzite lattice parameters (a- and c-axis), and estimated average grain size of TiAlN films grown with different N2 contents, [N2], in the Ar/N2 gas mixture. For comparison, the lattice parameters of w-AlN are also included. The error in the thickness and roughness value is around 10%. For the other parameters, the error in the last significant digit is indicated in brackets.
a
[N2] (%)
Th (nm)
Ti (at.%)
Al (at.%)
N (at.%)
Ti/Al
r (nm)
a (nm)
c (nm)
Grain size (nm)
– 10 25 50 75 50
– 402 281 170 120 1250
– 19(1) 11(1) 11(1) 12(1) –
50 46(1) 38(1) 36(1) 34(1) –
50 35(3) 51(3) 53(3) 56(3) –
0 0.40(2) 0.30(2) 0.32(3) 0.35(3) –
– 0.20 0.50 0.38 0.63 1.46
0.311a – 0.322(4) 0.322(4) 0.320(4) 0.319(4)
0.498a – 0.499(2) 0.510(2) 0.523(2) 0.506(2)
– – 4.9(1) 4.0(1) 3.8(1) 6.2(3)
Data extracted from Ref. [32].
correlates with a progressive increase in the Ti/Al ratio for [N2] P 25% (see Table 1). The surface morphology of the TiAlN films is shown in Fig. 2, as imaged by AFM. The film surfaces are very smooth with root-mean square roughness values, r, below 2 nm (even for the sample grown for 6 h). The morphology of the sample grown with [N2] = 10% is very smooth and featureless (Fig. 2a), with r close to the bare substrate (r 0.15–0.20 nm). This growth condition can be considered as ultrasmooth [29]. In contrast, for [N2] = 25% (Fig. 2b) the surface displays a clear granular structure in the 20 nm range (note that the value is overestimated due to tip convolution effects). The samples grown with [N2] = 50% (Fig. 2c) and 75% (Fig. 2d) are smoother and rougher, respectively, than that produced at 25%. A detailed inspection of Fig. 2b reveals quite featureless background morphology, akin to that observed in Fig. 2a, together with higher morphological structures scattered on top. In Fig. 2d, a similar granular structure as that in Fig. 2b is also observed. Typically, the surface of polycrystalline w-AlN [30] or c-TiN [31] films produced by sputter growth display a much rougher surface with granular or pyramidal motifs as a result of the (projected) columnar microstructure. In the present case, the morphology is rather smooth, which suggests that the films are nanocrystalline (with very fine grains) or even amorphous. In the following section, the morphologies displayed in Fig. 2 are correlated with the crystal structure extracted from XRD. It should be noted that, for such nanograin sizes, tip convolution effects hamper a reliable comparison between the values derived from AFM images and the nanocrystal sizes obtained by XRD. Fig. 1. (a) ERDA-BIC spectrum and (b) compositional profile from a TiAlN film produced by reactive magnetron sputtering with [N2] = 25%.
3.2. Phase identification by X-ray and electron diffraction
sensitive to the element concentration regardless of its chemical state. In all cases, Al-rich films are obtained under our working conditions, despite of the Ti/Al = 1 ratio of the target. In the case of TiAlN films grown by sputtering of composite targets, higher Al contents than the nominal value of the target are typically obtained [27,28]. This fact has been explained in the referred works by different sputtering rates of the two nitrides (higher for AlN than for TiN) and poisoning states of the Ti and Al atoms in the target, together with scattering and angular losses of the sputter flux. In this line, note that the Ti/Al ratio increases with [N2] for contents equal or above 25% (where stoichiometric nitrides are formed), i.e. as the compound mode operation of the target is enhanced. The element concentration shown in Fig. 1b is expressed in atoms/cm2 (projected or areal density), which can be transformed into a depth scale by assuming a realistic film density. Inversely, it can be used to calculate the density if the film thickness is known, as in our case. The thus obtained mass density values increase from 3.6 to 4.2 g/cm3 with [N2], which lie in between those for c-TiN (5.21 g/cm3) and w-AlN (3.26 g/cm3). This trend
The structure evolution of TiAlN films as a function of the [N2] is followed in Fig. 3 by XRD. The reflections from w-AlN and cubic binary structures (c-TiN and c-AlN) are also indicated in the upper panel. Note that the cubic binary structure shrinks (higher 2h angles) from c-TiN to c-AlN due to the smaller Al than Ti atoms. Due to the relatively low thickness, some features from the Si(1 0 0) substrate are observed around 55°. In correlation with the formation of a sub-stoichiometric nitride at [N2] = 10% (Table 1), the crystal structure of this sample is rather poor and only a broad and weak reflection is obtained around 2h = 38°. This broad feature may originate from highly-disordered cubic nitride structures (note the relatively high Ti content) or metal clusters that have not reacted with N. However, this assignment is ambiguous due to the superposition of several reflections in the same region. The disordered structure of this film also correlates with the smooth and featureless morphology shown in Fig. 2a. The diffraction pattern for samples grown with [N2] P 25% indicate the formation of a hexagonal phase with wurzite structure. The (0 0 2) and (1 0 3) reflections are mainly observed, although additional weak signals from (1 0 0), (1 0 2), to (1 1 0) reflections can also be identified at 2h of 32°, 49°, and 58°, respectively,
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Fig. 2. AFM images (750 750 nm2) of TiAlN films grown for 60 min with [N2] = 10% (a), 25% (b), 50% (c) and 75% (d).
Fig. 3. Glancing-angle XRD scans from TiAlN films grown with different N2 contents, [N2], in the gas mixture. The reflections from reference binary compounds of c-TiN (j), c-AlN (h) and w-AlN (4) are indicated in the upper panel. Some features from the Si(1 0 0) substrate are observed around 55°.
in the film grown with [N2] = 25%. The (0 0 2) and (1 0 3) positions are shifted to lower 2h values with respect to those of w-AlN indicating an expansion of the wurzite lattice. The reflections also appear at lower 2h by increasing [N2]. The corresponding lattice parameters derived from the positions of the (1 0 3) and (0 0 2) reflections are summarised in Table 1. The reference values from w-AlN are also included as extracted from Ref. [32]. For [N2] P 25%, the c-axis lattice parameter increases with the Ti/Al ratio whereas the a-axis remains more or less constant. Such a trend could be explained by the incorporation of Ti within the hexagonal network and/or the emergence of strain. It should be noted that Holec et al. [33] predicted that the a-axis lattice parameter should decrease while the c-axis remains constant with the AlN mole fraction. This opposite behaviour is an additional indication of the eventual significance of stress in our samples. In fact, samples grown with [N2] P 25% show poor adhesion to the Si substrate for large thickness (>1 lm).
The sample grown with [N2] = 25% shows the highest crystalline quality since it shows the larger number of reflections from the wurzite phase and, in the case of (0 0 2) and (1 0 3) peaks, they are narrower and more intense than for the other samples (this result may be influenced by the larger thickness of this film). The width of the reflections indicate the formation of very fine nanocrystalline films with a mean domain size of just a few nm’s, as derived from the Scherrer formula [34] (note that the obtained value may be overestimated by other broadening effects such as strain). This crystalline improvement correlates with the observation of a neat surface nanogranular surface structure by AFM in Fig. 2b. In order to meaningfully compare the data obtained by AFM (roughness) and XRD (nanocrystal size), the different thickness of the films should be taken into account. Thus, the roughness decrease for [N2] P 50% could be due, to some extent, to the parallel diminution in the film thickness [35]. Likewise, the reduction in grain size (Table 1) could be also correlated to the film thickness reduction. However, for the highest [N2] (thinnest film), the roughness is higher. This fact suggests that the roughness and grain size dependences on growth conditions could have two different sources, the thickness and the crystallinity. The microstructure evolution of TiAlN as a function of the film thickness has been studied by growing a sample for 360 min at [N2] = 50%. As shown in Table 1, the thickness of this sample is around 1250 nm. As expected [35], the roughness increases with the film thickness. Fig. 4a compares the glancing-angle XRD pattern for this sample with that of the thinner film (170 nm) grown under equivalent conditions (same curve as in Fig. 3). A similar structure is found in both samples, revealing a dominant hexagonal phase with lattice parameters larger than those of w-AlN. Also, a small contribution from a ternary c-Ti1xAlxN phase is observed (as derived by the lattice parameters in between those of c-TiN and c-AlN) for the film with the larger thickness. The grain size (Table 1) is larger for the thicker film, as expected from the grain growth with time and the columnar microstructure. The latter can be seen from the cross-sectional bright-field TEM image in Fig. 4b. Fig. 4c shows the selected area electron diffraction (SAED) pattern from the circular region within the TiAlN film indicated with a dashed circle in Fig. 4b. The rings in the SAED pattern are in agreement with the observations by XRD and can be explained by an expansion of the wurzite lattice with respect to w-AlN of around 2.5% (solid rings). The analysis also confirms the presence of a small amount of a cubic phase (dashed rings) with a 2.5% compression with respect to c-TiN (coming from the cubic ternary solid solution). Finally, the hexagonal phase presents a (0 0 2) texture as the corresponding diffraction ring (d002 = 0.261 nm) has higher intensity in the direction parallel to the substrate normal (vertical direction in Fig. 4c).
3.3. X-ray Absorption Near-Edge Structure (XANES) In the previous section, the formation of a hexagonal phase with larger lattice parameters than w-AlN (2.5%) has been identified. This result has been discussed in terms of the formation of hexagonal ternary phases with different Ti/Al ratios and/or strain. To verify the formation of a ternary phase, the TiAlN samples grown for 1 h with different [N2] values in the gas mixture have been analysed by XANES at the Ti2p, Al1s, and N1s edges. Although the XANES can also be slightly distorted by the presence of strain [36,37], the overall lineshape is not modified since the local coordination is not altered. Fig. 5 shows the spectra of TiAlN films together with the reference spectra from binary c-TiN and w-AlN. The Al1s edge spectrum from a c-Ti1xAlxN solid solution with x = 0.35 is also included as reference of Al arrangements within the cubic lattice [17].
R. Gago et al. / Journal of Alloys and Compounds 561 (2013) 87–94
Fig. 4. (a) Glancing incidence XRD patterns from 170 and 1250 nm thick TiAlN films grown with [N2] = 50%. The reflections from c-TiN (j), c-AlN (h) and w-AlN (4) are indicated. (b) Cross-sectional TEM image and (c) associated SAED pattern from the thicker TiAlN film. The position of the selected area aperture is marked by a dashed circle in (b). The rings in the SAED pattern have been assigned to hexagonal (solid) and cubic (dashed) nitride phases with larger and shorter lattice spacing, respectively, than the binary compounds. Note that several rings do overlap (dash-dot circles) for both structures.
The Ti2p spectra of TiAlN show four main contributions labelled as A, B, C, and D. This edge originates from 2p to 3d transitions and, hence, shows a L2 (C and D features) and L3 (A and B features) complex structure due to the spin–orbit interaction of the Ti-2p levels. The 3d levels in transition metals can also be split due to crystal-field effects [38]. Typically, the L2,3 splitting in the Ti2p XANES spectrum of c-TiN is 6 eV whereas the splitting due to crystal-field effects is around 1.5 eV [39]. In the case of covalent bonding, such as in TiO2, the L2,3 edges present a well-defined and strong crystal-field splitting (3 eV), with distinct spectral features according to the different coordination geometries [40].
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However, this effect is less pronounced in the case of materials with metallic character [41], such as the case of pure Ti or c-TiN, presenting broader features. Accordingly, the Ti2p XANES spectrum from c-TiN in Fig. 5 shows two main broad features around 459 and 465 eV, in agreement with the results from Soriano et al. [39]. There is also a small contribution of Ti–O bonds from surface oxidation as derived by the peak at 459.5 eV [39]. Soriano et al. [39] has shown that the Ti–O feature is promoted upon progressive surface oxidation until the spectrum resembles that of TiO2. Note that the contribution from Ti–O is absent in the TiAlN films produced with a compound target, related to the higher oxidation resistance of Ti atoms in the ternary nitride as compared to cTiN. Remarkably, the position and relative intensity of the fine structure of the L2,3 edges in TiAlN differ from that of c-TiN. The presence of fine structure suggests a higher contribution from crystal-field effects and points out toward a more covalent bonding configuration around Ti atoms. Interestingly, the relative intensity of features A and B is inversed for [N2] P 25%. The different spectral features for [N2] P 25% should be correlated with the formation of a nanocrystalline hexagonal phase (see Fig. 3). Regarding the Al1s edge, the w-AlN reference sample shows a similar spectrum as that reported in the literature [42,43]. The fine structure in the spectrum is related to transitions from Al-1s to p and d hybridized orbitals, showing a low-energy peak around 1565 eV and a more intense and broader feature around 1571 eV. As predicted theoretically [46,44], the Al1s edge in AlN with rock-salt coordination lacks the low-energy peak characteristic of w-AlN and the more intense feature is downshifted 1 eV with respect to the maxima in the w-AlN spectrum, as indeed observed experimentally [20,21]. Based on this, Al sites with cubic arrangement have been identified in cubic ternary solid solutions [15,17], as highlighted by the reference spectrum for c-Ti1xAlxN (x 0.35) in Fig. 5. In line with the composition and structural analysis, the amorphous and sub-stoichiometric character of the sample grown with [N2] = 10% results in broad features whereas TiAlN samples grown with [N2] P 25% show similar Al1s spectra. In the latter case, the spectral lineshape resembles that of w-AlN as extracted by the presence of a low-energy peak (feature E) followed by a more intense one at higher energies (feature F). Note that also a hump close to 1582 eV is present in the spectra of both TiAlN samples and w-AlN (feature G). The latter feature has also been predicted in ternary wurzite phases and, moreover, its position should not be influenced by the Ti/Al ratio [12] as observed here. In agreement with XRD, the sample grown with [N2] = 25% shows narrower E and F features due to the higher structural ordering (see Table 1). Remarkably, features E and F of TiAlN are shifted to higher energies with respect to w-AlN. This suggests that, although Al sites are in hexagonal arrangements they slightly differ from those at w-AlN. In particular, the peak shifts could result from changes in the wurzite lattice parameters (including bond angle and distance distortions) and/or the bonding character [12] with Ti incorporation. Hence, in conjunction with the changes in the Ti2p spectra, the Al1s edge provides another indirect proof of the formation of a ternary hexagonal phase. The N1s spectra can also yield relevant information about the formation of a ternary compound and deserves special attention, since the N1s edge of ternary cubic and hexagonal nitrides have been previously studied theoretically [12]. On the one hand, the spectrum from c-TiN can be related to transitions from N-1s to unoccupied N-2p states with two clear regions related to the hybridization with the empty Ti-3d (397–404 eV) and Ti-4sp (404–414 eV) bands [45]. The first region shows two components due to the splitting of the Ti-3d states into t2g and eg sub-bands due to crystal-field effects [38]. On the other hand, the lineshape of the N1s spectra in w-AlN corresponds to transitions from N-1s to unoccupied N-2p states [46]. The spectrum shows a clear
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Fig. 5. Ti2p, Al1s, and N1s XANES from TiAlN films grown with different [N2]. The reference spectra from c-TiN and w-AlN compounds grown under similar conditions are displayed too. The Al1s spectrum from a ternary c-Ti1xAlxN solid solution (x = 35) is also included as reference for Al sites in the cubic structure. The spectral intensity has been normalised to one in all cases and the spectra are shifted vertically for clarity.
triple-peak structure whose relative intensity strongly depends on the X-ray incidence angle [47]. The overlap of the N1s features from the reference binary nitrides adds some complexity to the identification between cubic and hexagonal environments in ternary compounds. However, the filling of states around 405 eV (features I–K) in the spectra of TiAlN films indicates the relevant contribution from N sites surrounded by Al atoms in wurzite-like arrangements. The fine structure of the spectrum from w-AlN within 400–410 eV is not present in the TiAlN films. However, it has been predicted [12] that such triple-peak structure is levelled out in wurzite ternary compounds. As in the Al1s edge, the N1s features (I–K) are shifted to higher energies with respect to the positions in w-AlN, which supports the formation of a ternary wurzite phase. There is also a significant contribution from feature H, with energy close to the low-energy peak in the N1s spectrum of c-TiN. This might be interpreted as a certain fraction of N atoms in octahedral coordination. A similar observation was reported by Gîrleanu et al. [48] for Ti0.14Al0.86N films studied by ELNES and attributed to cubic arrangements at grain boundaries. However, in a previous publication [15], the same authors found that a significant fraction of Ti atoms are incorporated in the fourfold sites of the wurzite lattice. Also, the spectrum reported in [48] has different fingerprints as those presented here and, in their case, the triple-peak structure characteristic of w-AlN was clearly resolved. In order to study the N arrangements in more detail, we have first considered the hypothesis that the TiAlN films consist of a mixture of binary phases (that is, excluding the formation of ternary phases). Under such assumption, the N1s spectra of TiAlN in Fig. 5 should be a linear combination of that from c-TiN and wAlN. In order to verify this, the c-TiN spectrum has been subtracted from that of TiAlN and the result is shown in Fig. 6. The subtracted spectrum (solid line) has some similarities with w-AlN (dashed line) but important features are missing, specially the triple-peak structure between 400 and 410 eV. Therefore, we cannot interpret the spectrum of TiAlN films as a sole combination of binary phases. This also agrees with the novel features described for the Ti2p and Al1s edges. Fig. 7 compares our spectra with those predicted for w-Ti1xAlxN phases by Holec et al. [12]. Reproducing the experimental spectra of known structures such as c-TiN and w-AlN is still current
state-of-art research and, obviously, we do not expect to match our experimental spectra where unknown and/or mixed phases may be present. Rather, we are looking for qualitative agreements. Moreover, additional broadening should be applied to the calculations in order to approximate to the experimental conditions [12]. For our comparison, the first peak of the experimental spectrum of w-AlN has been used for energy calibration. First, as a test of the comparison, we can see that there is a good correlation between the measured spectrum of w-AlN and the calculations. Second, the subtracted spectrum from Fig. 6 is found to be similar to that of a w-Ti1xAlxN structure with x = 0.875. This could be interpreted as the result from a mixture between a hexagonal ternary phase with a lower Al content than the nominal one (x = 0.75) and c-TiN. However, it should be noted that no cubic phase has been detected by XRD (Fig. 3) in these samples grown for 1 h (those analysed by XANES). This would imply that the cubic arrangements, if present, are located at the grain boundaries [16]. Taking this into account, the contribution from grain boundaries (and, correspondingly, the c-TiN signal) should increase with [N2] since the crystal-
Fig. 6. N1s spectrum for the TiAlN sample grown with [N2] = 25% after subtraction (solid line) of the c-TiN spectrum and comparison with w-AlN (dashed line).
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in the films. This trend suggests the incorporation of Ti in the hexagonal network, which causes expansion of the lattice. The study of the Ti2p, Al1s, and N1s XANES spectra confirm the formation of a ternary hexagonal compound since the spectra cannot be fitted with a linear combination of the spectra for the binary nitrides and the spectral features are shifted with respect to those of wAlN. Moreover, the previous assumption is supported by comparison with theoretical modelling reported in the literature. Acknowledgments
Fig. 7. Comparison of experimental (dotted curves) N1s spectra of w-AlN and Ti0.25Al0.75N films with calculated edges (thin curves) of w-Ti1xAlxN phases with x = 1 and 0.75, respectively. The subtracted spectrum from Fig. 5 is also compared with a w-Ti1xAlxN phase with x = 0.875. The theoretical curves have been adapted from Ref. [12].
line quality deteriorates (see Fig. 3 and Table 1). However, the N1s spectra in Fig. 5 for TiAlN films grown with [N2] P 25% show a similar relative intensity of feature H (taken as a measure of the c-TiN content in the films). Remarkably, the predicted N1s edge from wTi1xAlxN with x = 0.75 can also be satisfactorily correlated with the fine structure of the experimental N1s XANES spectrum of the Ti0.25Al0.75N film, even including the low energy peak around 398.5 eV (feature H in Fig. 5). Based on the calculations reported in Ref. [12] for Al-rich TiAlN, it would be expected that N sites near Al tend to have a four-coordination (wurzite-like) whereas fivecoordination would be preferred in the vicinity of Ti neighbours. Therefore, we attribute feature H in the N1s spectra of TiAlN films not to cubic arrangements but rather to N atoms surrounded by Ti atoms in hexagonal arrangements. In fact, the N1s spectrum from hexagonal TiN (BN structure) is predicted to have similar features as in the case of c-TiN [12]. It should be mentioned that the calculations in Ref. [12] were limited to the N1s edge and, therefore, it would be desirable to complete the study of the other elemental edges to attain more solid conclusions. Hence, the present experimental results support previous theoretical modelling and could be implemented in further theoretical modelling to understand the formation and electronic structure of ternary nitride phases.
4. Conclusions TiAlN thin films have been grown by reactive DC magnetron sputtering in an Ar/N2 atmosphere from a Ti50Al50 compound target. ERDA-BIC shows that the film composition is Al-rich and that stoichiometric nitrides are produced above a certain threshold of N2 contents in the gas mixture ([N2] P 25%). AFM shows that the surface morphology of the TiAlN films is very smooth (r > 2 nm). In the case of stoichiometric nitrides, glancing-incidence XRD shows that polycrystalline films with wurzite structure are obtained, the c-axis lattice parameter increasing with the Ti/Al ratio
We thank G.S. Fox-Rabinovich for providing cemented carbide inserts and T. Regier for beamline assistance at SGM (CLS). This work is part of the research carried out within the bilateral 2010RU0025 project between CSIC (Spain) and RFBR (Russia). Additional financial support from grants FIS2012-38866-C05-05 (MECO, Spain), AVANSENS no. S2009/PPQ-1642 (CAM, Spain), and CSD2008-00023 (MICINN, Spain) is acknowledged. ERD-BIC measurements have been supported by the EC Integrating Activity ‘‘Support of Public and Industrial Research Using Ion Beam Technology (SPIRIT)’’ under contract no. 227012. The synchrotron work at CLS (Canada) has been supported by NSERC, NRC, CIHR and the University of Saskatchewan. F.S. thanks the EFRE Funds of the European Commission for support of activities within the AME-Lab project. Finally, ARC acknowledges funding from SFRH/BPD/74095/2010 and PTDC/CTM/100756/2008 grants (FTC, Portugal). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
X-ray absorption near-edge structure of hexagonal ternary phases in sputter-deposited TiAlN films R. Gago a,⇑, F. Soldera b, R. Hübner c, J. Lehmann c, F. Munnik c, L. Vázquez a, A. Redondo-Cubero d, J.L. Endrino a,e a
Instituto de Ciencia de Materiales de Madrid, Consejo Superior de Investigaciones Científicas, E-28049 Madrid, Spain Department of Materials Science & Engineering, Saarland University, D-66123 Saarbruecken, Germany Institute of Ion Beam Physics and Materials Research, Helmholtz-Zentrum Dresden-Rossendorf, D-01314 Dresden, Germany d Instituto Tecnológico e Nuclear, Instituto Superior Técnico, Universidade Técnica de Lisboa, 2686-953 Sacavém, Portugal e Abengoa Research S.L., c/Energía Solar 1, Palmas Altas, E-41014 Seville, Spain b c
a r t i c l e
i n f o
Article history: Received 12 December 2012 Received in revised form 21 January 2013 Accepted 23 January 2013 Available online 13 February 2013 Keywords: Nitride materials Vapour deposition Atomic scale structure NEXAFS/XANES
a b s t r a c t Titanium aluminium nitride (TiAlN) coatings have been grown by reactive (Ar/N2) direct-current magnetron sputtering from a Ti50Al50 compound target. The film composition has been quantified by ion beam analysis showing the formation of Al-rich nitrides (Ti/Al 0.3), with stoichiometric films for N2 contents in the gas mixture equal or above 25%. The surface morphology of the films has been imaged by atomic force microscopy, showing very smooth surfaces with roughness values below 2 nm. X-ray and electron diffraction patterns reveal that the films are nanocrystalline with a wurzite (w) structure of lattice parameters larger (2.5%) than those for w-AlN. The lattice expansion correlates with the Ti/Al ratio in stoichiometric films, which suggests the incorporation of Ti into w-AlN. The atomic environments around Ti, Al and N sites have been extracted from the X-ray absorption near-edge structure (XANES) by recording the Ti2p, Al1s and N1s edges, respectively. The analysis of the XANES spectral lineshape and comparison with reported theoretical calculations confirm the formation of a ternary hexagonal phase. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Transition metal nitride materials have been commercially available for many decades as hard, wear-resistant, and decorative coatings, as well as electrical contacts and conductive diffusion barriers in electronic devices [1]. Among them, titanium nitride (TiN) is widely used as protective coating material in cutting tools, mechanical components and medical implants due to its high hardness, non-toxicity, low friction, and chemical inertness [2]. Ternary compounds within the Ti–Al–N system have been extensively studied as an alternative to TiN overcoats for high-temperature applications since Al improves the wear and oxidation resistance, as well as the thermal stability of the coated system [3]. In the case of ternary alloys, a crucial aspect to tune the final properties is an accurate control of both the composition and structure of the thin films. Hence, the production of ternary TiAlN compounds has been mainly focused on increasing the Al content in the cubic TiN (c-TiN) structure (solid solution) to maximize oxidation resistance while preventing the segregation of binary phases [4]. However, c-TiN can only accommodate up to a certain amount of AlN molar fractions (x) ⇑ Corresponding author at: Instituto de Ciencia de Materiales de Madrid, Consejo Superior de Investigaciones Científicas, E-28049 Madrid, Spain. Tel.: +34913349090. E-mail address: [email protected] (R. Gago). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.01.130
within a metastable N-saturated solid solution (c-Ti1xAlxN) [5]. As the solubility limit is exceeded, the film structure undergoes a spinodal decomposition (i.e., c-Ti1xAlxN domains with different composition) [6] or phase segregation of c-Ti1xAlxN and (hexagonal) wurzite aluminium nitride (w-AlN) [7]. Theoretically, the critical AlN molar fraction is predicted to be x 0.7 [8]. Experimentally, a broad range of solubility limits have been published (0.4 < x < 0.9) [9], which could be understood based on the multiple growth conditions employed. Contrary to c-Ti1xAlxN, the formation of hexagonal ternary single-phases of w-Ti1xAlxN (that is, the incorporation of Ti into w-AlN) has not been studied in detail so far. Such phases are softer than the cubic counterpart but present an excellent oxidation resistance [10]. Hence, they can be suitable candidates for certain technological applications. Further, the study of w-Ti1xAlxN phases could be relevant to understand the formation and electronic structure of nanolaminated ternary materials in the Ti–Al– N system with hexagonal structure [11]. In this case, it has been predicted [12] that the wurzite structure can only accommodate Ti contents up to 20 at.% (x 0.6), being unstable above this threshold. X-ray diffraction (XRD) is the most common technique used to study the microstructure of Ti–Al–N coatings. However, XRD probes only the crystallized domains and, hence, information from the eventual participation of amorphous nitrides or grain boundaries
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is lacking [13]. Further, the position of the Bragg reflections can overlap for different structures or be affected by other issues such as lattice strain. Complementary, the near-edge structure (NES) extracted from electron energy loss (ELNES) or X-ray absorption (XANES) measurements can be used to solve complex systems and discern the coexistence of different phases, even in amorphous or disordered structures. Typically, XANES presents a better spectral resolution than ELNES (0.1 eV instead of 1 eV) although it requires monochromatic X-rays from a synchrotron facility [14]. These techniques probe the density of unoccupied states and yield local order information with element selectivity by recording specific edges. In this way, the simultaneous analysis of the characteristic absorption edge(s) for each element can be unequivocally used for phase identification (distinct spectral features are used as phase fingerprints). Despite the potential of this analysis, there are only few XANES reports addressing the formation [15–17] and thermal stability [18] of ternary TiAlN compounds. It should be noted that c-AlN is a metastable phase and it is only produced under highpressure [19], whereas TiN is mechanically unstable in the fourcoordinated wurzite structure [12]. Therefore, identification of cubic or hexagonal ternary environments by XANES may be complex due to the lack of reference nitrides with Al and Ti sites occupied in the rock-salt or hexagonal structure, respectively. Nevertheless, XANES spectra of c-AlN have been successfully recorded by stabilizing this phase in short-period TiN/AlN multilayers under compressive stress [20] and upon segregation in low carbon aluminiumkilled steels [21]. Alternatively, theoretical modelling of ELNES or XANES spectra can be used to predict and understand the bonding behaviour of ternary TiAlN phases with either hexagonal or cubic structure [12]. However, experimental evidence to support such predictions is still needed, especially, in the case of hexagonal phases. In this work, the formation of hexagonal TiAlN ternary phases is identified by the combination of XRD and XANES. Soft X-rays are used to study the atomic arrangements around Ti, Al and N by recording the Ti2p, Al1s and N1s XANES spectra, respectively. The presence of distinct spectral features in the XANES measurements with respect to binary nitrides and comparison with reported theoretical results support the formation of a ternary phase. To the best of our knowledge, XANES of w-Ti1xAlxN phases has not been reported experimentally so far. This information will be relevant to identify the formation of ternary phases in future work and will provide experimental data for further modelling and understanding the electronic structure of ternary nitride systems. 2. Materials and methods 2.1. Sample preparation TiAlN films were grown on Si(1 0 0) substrates by reactive direct-current (DC) magnetron sputtering from a 3’’ Ti50Al50 (in at.%) composite target. The cathode was facing the grounded substrates at a working distance of 15 cm. The base and working pressure were 104 and 0.2 Pa, respectively. Different Ar (99.9995% purity grade) and N2 (99.9995% purity grade) contents in the gas mixture were used to influence the film properties. For this purpose, the individual Ar (UAr) and N2 (UN2) flows were adjusted while keeping the total flow (pressure) constant at 20 standard cubic centimeter per minute (sccm). In this way, the growth has been parameterized as a function of the N2 content in the gas mixture, [N2] = UN2/ (UAr + UN2). For plasma generation, a DC voltage with an overall power of 200 W was applied to the cathode. The depositions were carried out for 1 h at room temperature. Additionally, a thick sample (>1 lm) was grown for 6 h with [N2] = 50% to study the microstructure evolution with thickness. In this case, the film was grown onto a cemented carbide insert to improve the adhesion with respect to Si(1 0 0) substrates. 2.2. Sample characterisation The thickness of the TiAlN films was determined with a Dektak-150 (VeecoÒ) mechanical profiler by measuring the step height created by partially masking the substrate during growth. The composition of the layers was determined by
heavy-ion elastic recoil detection analysis (ERDA) with a 35 MeV Cl7+ beam impinging at 15° with respect to the sample surface. The scattered ions and recoils were detected with a Bragg ionisation chamber (BIC) located at a scattering angle of 31° that discriminates the detected particle according to their atomic number (Z). The measured energy spectrum of each element was converted into compositional depth profiles by means of the analysis program WiNDF [22] using the stopping power data from Ziegler et al. [23]. The resulting surface morphology was imaged by atomic force microscopy (AFM) operating in the dynamic mode with an Agilent PicoPlus 5500. AFM probes with a nominal radius of 8 nm and opening angle smaller than 52° were used. Image editing and analysis was done with the Gwyddion freeware package [24]. The microstructure of the samples was examined by glancing-angle XRD measurements using a D5000 (BRUKER AXS) diffractometer with Cu Ka radiation (k = 1.5418 Å). The data were collected with an incidence angle of 1° within the scattering range, 2h, of 30–80°. Complementary, the bonding structure with element sensitivity to Ti, Al and N sites was studied by XANES with soft X-rays at the Canadian Light Source (CLS) synchrotron using the SGM beamline, which provides an energy resolution of 0.1 eV. The measurements were performed in the total electron yield (TEY) mode with the sample normal positioned parallel to the X-ray beam. Finally, the microstructural evolution was examined by crosssectional transmission electron microscopy (TEM) with an image-corrected TITAN 80-300 (FEI) microscope operated at 300 kV. For this purpose, a slice from the thick TiAlN film was prepared in a focused ion beam and scanning electron microscope dual beam system (FEI Helios 600) by using the ‘‘in situ’’ lift-out technique. The foil was thinned to electron transparency, first, with an acceleration voltage of 30 kV and, later, with 5 kV (for at least 2 min in each side) to minimize any possible preparation artifacts.
3. Results and discussion 3.1. Thickness, composition and surface morphology The thickness of the TiAlN films as a function of N2 content in the gas mixture, [N2], is shown in Table 1. Clearly, the thickness decreases progressively with [N2] or, equivalently, the deposition rate drops from 7 to 2 nm/min. This trend is the result of target poisoning (nitride formation) and the transition toward the sputtering compound mode by increasing the reactive gas content [25]. As an example, the ERDA-BIC data from a TiAlN film grown with [N2] = 25% is shown in Fig. 1a. The signals from the elements in the film and the scattered Cl projectiles have been labelled in the graph. The potential of this analysis is that well-resolved signals are obtained, specially, for light elements. In addition, heavier elements such as Ti provide double information from Ti recoils and Cl projectiles scattered from Ti target atoms. For each detected particle, lower energies imply scattering events deeper in the sample. The energy onset for elements at the surface is given by the kinematic factor of the collision (this onset is higher for heavier elements). Note that oxygen is only detected at the surface as an indication of surface contamination and/or oxidation, and at the substrate/film interface from the native oxide of the Si(1 0 0) substrate. In addition, a negligible concentration of C is detected, showing the high purity grade of the process. In the ERDA-BIC map, the signal from each element can be projected independently to the energy (abscissa) axis and, in this way, the element depth profile can be derived. As an example, the compositional profile derived from the ERDA-BIC data of Fig. 1a is plotted in panel b. It can be seen that a stoichiometric nitride is obtained with a composition close to Ti0.25Al0.75N. The average film composition has been quantified by integrating the compositional profiles derived from ERDA-BIC over the whole film thickness. The resulting values are included in Table 1 showing an increase of the N content with [N2]. A sub-stoichiometric nitride is obtained for the lowest value of [N2] = 10%, which means that part of the metal atoms do not react to form a nitride due to deficient N supply. On the contrary, stoichiometric nitrides (N = 50 at.%) are formed for [N2] P 25%, with slightly over-stoichiometric values for [N2] P 50%. In this case, the slight excess of N is probably incorporated in the form of embedded N2 molecules, as observed for the growth of TiN films [26]. Note that ERDA-BIC is
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Table 1 Film thickness (Th), composition, surface roughness (r), wurzite lattice parameters (a- and c-axis), and estimated average grain size of TiAlN films grown with different N2 contents, [N2], in the Ar/N2 gas mixture. For comparison, the lattice parameters of w-AlN are also included. The error in the thickness and roughness value is around 10%. For the other parameters, the error in the last significant digit is indicated in brackets.
a
[N2] (%)
Th (nm)
Ti (at.%)
Al (at.%)
N (at.%)
Ti/Al
r (nm)
a (nm)
c (nm)
Grain size (nm)
– 10 25 50 75 50
– 402 281 170 120 1250
– 19(1) 11(1) 11(1) 12(1) –
50 46(1) 38(1) 36(1) 34(1) –
50 35(3) 51(3) 53(3) 56(3) –
0 0.40(2) 0.30(2) 0.32(3) 0.35(3) –
– 0.20 0.50 0.38 0.63 1.46
0.311a – 0.322(4) 0.322(4) 0.320(4) 0.319(4)
0.498a – 0.499(2) 0.510(2) 0.523(2) 0.506(2)
– – 4.9(1) 4.0(1) 3.8(1) 6.2(3)
Data extracted from Ref. [32].
correlates with a progressive increase in the Ti/Al ratio for [N2] P 25% (see Table 1). The surface morphology of the TiAlN films is shown in Fig. 2, as imaged by AFM. The film surfaces are very smooth with root-mean square roughness values, r, below 2 nm (even for the sample grown for 6 h). The morphology of the sample grown with [N2] = 10% is very smooth and featureless (Fig. 2a), with r close to the bare substrate (r 0.15–0.20 nm). This growth condition can be considered as ultrasmooth [29]. In contrast, for [N2] = 25% (Fig. 2b) the surface displays a clear granular structure in the 20 nm range (note that the value is overestimated due to tip convolution effects). The samples grown with [N2] = 50% (Fig. 2c) and 75% (Fig. 2d) are smoother and rougher, respectively, than that produced at 25%. A detailed inspection of Fig. 2b reveals quite featureless background morphology, akin to that observed in Fig. 2a, together with higher morphological structures scattered on top. In Fig. 2d, a similar granular structure as that in Fig. 2b is also observed. Typically, the surface of polycrystalline w-AlN [30] or c-TiN [31] films produced by sputter growth display a much rougher surface with granular or pyramidal motifs as a result of the (projected) columnar microstructure. In the present case, the morphology is rather smooth, which suggests that the films are nanocrystalline (with very fine grains) or even amorphous. In the following section, the morphologies displayed in Fig. 2 are correlated with the crystal structure extracted from XRD. It should be noted that, for such nanograin sizes, tip convolution effects hamper a reliable comparison between the values derived from AFM images and the nanocrystal sizes obtained by XRD. Fig. 1. (a) ERDA-BIC spectrum and (b) compositional profile from a TiAlN film produced by reactive magnetron sputtering with [N2] = 25%.
3.2. Phase identification by X-ray and electron diffraction
sensitive to the element concentration regardless of its chemical state. In all cases, Al-rich films are obtained under our working conditions, despite of the Ti/Al = 1 ratio of the target. In the case of TiAlN films grown by sputtering of composite targets, higher Al contents than the nominal value of the target are typically obtained [27,28]. This fact has been explained in the referred works by different sputtering rates of the two nitrides (higher for AlN than for TiN) and poisoning states of the Ti and Al atoms in the target, together with scattering and angular losses of the sputter flux. In this line, note that the Ti/Al ratio increases with [N2] for contents equal or above 25% (where stoichiometric nitrides are formed), i.e. as the compound mode operation of the target is enhanced. The element concentration shown in Fig. 1b is expressed in atoms/cm2 (projected or areal density), which can be transformed into a depth scale by assuming a realistic film density. Inversely, it can be used to calculate the density if the film thickness is known, as in our case. The thus obtained mass density values increase from 3.6 to 4.2 g/cm3 with [N2], which lie in between those for c-TiN (5.21 g/cm3) and w-AlN (3.26 g/cm3). This trend
The structure evolution of TiAlN films as a function of the [N2] is followed in Fig. 3 by XRD. The reflections from w-AlN and cubic binary structures (c-TiN and c-AlN) are also indicated in the upper panel. Note that the cubic binary structure shrinks (higher 2h angles) from c-TiN to c-AlN due to the smaller Al than Ti atoms. Due to the relatively low thickness, some features from the Si(1 0 0) substrate are observed around 55°. In correlation with the formation of a sub-stoichiometric nitride at [N2] = 10% (Table 1), the crystal structure of this sample is rather poor and only a broad and weak reflection is obtained around 2h = 38°. This broad feature may originate from highly-disordered cubic nitride structures (note the relatively high Ti content) or metal clusters that have not reacted with N. However, this assignment is ambiguous due to the superposition of several reflections in the same region. The disordered structure of this film also correlates with the smooth and featureless morphology shown in Fig. 2a. The diffraction pattern for samples grown with [N2] P 25% indicate the formation of a hexagonal phase with wurzite structure. The (0 0 2) and (1 0 3) reflections are mainly observed, although additional weak signals from (1 0 0), (1 0 2), to (1 1 0) reflections can also be identified at 2h of 32°, 49°, and 58°, respectively,
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Fig. 2. AFM images (750 750 nm2) of TiAlN films grown for 60 min with [N2] = 10% (a), 25% (b), 50% (c) and 75% (d).
Fig. 3. Glancing-angle XRD scans from TiAlN films grown with different N2 contents, [N2], in the gas mixture. The reflections from reference binary compounds of c-TiN (j), c-AlN (h) and w-AlN (4) are indicated in the upper panel. Some features from the Si(1 0 0) substrate are observed around 55°.
in the film grown with [N2] = 25%. The (0 0 2) and (1 0 3) positions are shifted to lower 2h values with respect to those of w-AlN indicating an expansion of the wurzite lattice. The reflections also appear at lower 2h by increasing [N2]. The corresponding lattice parameters derived from the positions of the (1 0 3) and (0 0 2) reflections are summarised in Table 1. The reference values from w-AlN are also included as extracted from Ref. [32]. For [N2] P 25%, the c-axis lattice parameter increases with the Ti/Al ratio whereas the a-axis remains more or less constant. Such a trend could be explained by the incorporation of Ti within the hexagonal network and/or the emergence of strain. It should be noted that Holec et al. [33] predicted that the a-axis lattice parameter should decrease while the c-axis remains constant with the AlN mole fraction. This opposite behaviour is an additional indication of the eventual significance of stress in our samples. In fact, samples grown with [N2] P 25% show poor adhesion to the Si substrate for large thickness (>1 lm).
The sample grown with [N2] = 25% shows the highest crystalline quality since it shows the larger number of reflections from the wurzite phase and, in the case of (0 0 2) and (1 0 3) peaks, they are narrower and more intense than for the other samples (this result may be influenced by the larger thickness of this film). The width of the reflections indicate the formation of very fine nanocrystalline films with a mean domain size of just a few nm’s, as derived from the Scherrer formula [34] (note that the obtained value may be overestimated by other broadening effects such as strain). This crystalline improvement correlates with the observation of a neat surface nanogranular surface structure by AFM in Fig. 2b. In order to meaningfully compare the data obtained by AFM (roughness) and XRD (nanocrystal size), the different thickness of the films should be taken into account. Thus, the roughness decrease for [N2] P 50% could be due, to some extent, to the parallel diminution in the film thickness [35]. Likewise, the reduction in grain size (Table 1) could be also correlated to the film thickness reduction. However, for the highest [N2] (thinnest film), the roughness is higher. This fact suggests that the roughness and grain size dependences on growth conditions could have two different sources, the thickness and the crystallinity. The microstructure evolution of TiAlN as a function of the film thickness has been studied by growing a sample for 360 min at [N2] = 50%. As shown in Table 1, the thickness of this sample is around 1250 nm. As expected [35], the roughness increases with the film thickness. Fig. 4a compares the glancing-angle XRD pattern for this sample with that of the thinner film (170 nm) grown under equivalent conditions (same curve as in Fig. 3). A similar structure is found in both samples, revealing a dominant hexagonal phase with lattice parameters larger than those of w-AlN. Also, a small contribution from a ternary c-Ti1xAlxN phase is observed (as derived by the lattice parameters in between those of c-TiN and c-AlN) for the film with the larger thickness. The grain size (Table 1) is larger for the thicker film, as expected from the grain growth with time and the columnar microstructure. The latter can be seen from the cross-sectional bright-field TEM image in Fig. 4b. Fig. 4c shows the selected area electron diffraction (SAED) pattern from the circular region within the TiAlN film indicated with a dashed circle in Fig. 4b. The rings in the SAED pattern are in agreement with the observations by XRD and can be explained by an expansion of the wurzite lattice with respect to w-AlN of around 2.5% (solid rings). The analysis also confirms the presence of a small amount of a cubic phase (dashed rings) with a 2.5% compression with respect to c-TiN (coming from the cubic ternary solid solution). Finally, the hexagonal phase presents a (0 0 2) texture as the corresponding diffraction ring (d002 = 0.261 nm) has higher intensity in the direction parallel to the substrate normal (vertical direction in Fig. 4c).
3.3. X-ray Absorption Near-Edge Structure (XANES) In the previous section, the formation of a hexagonal phase with larger lattice parameters than w-AlN (2.5%) has been identified. This result has been discussed in terms of the formation of hexagonal ternary phases with different Ti/Al ratios and/or strain. To verify the formation of a ternary phase, the TiAlN samples grown for 1 h with different [N2] values in the gas mixture have been analysed by XANES at the Ti2p, Al1s, and N1s edges. Although the XANES can also be slightly distorted by the presence of strain [36,37], the overall lineshape is not modified since the local coordination is not altered. Fig. 5 shows the spectra of TiAlN films together with the reference spectra from binary c-TiN and w-AlN. The Al1s edge spectrum from a c-Ti1xAlxN solid solution with x = 0.35 is also included as reference of Al arrangements within the cubic lattice [17].
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Fig. 4. (a) Glancing incidence XRD patterns from 170 and 1250 nm thick TiAlN films grown with [N2] = 50%. The reflections from c-TiN (j), c-AlN (h) and w-AlN (4) are indicated. (b) Cross-sectional TEM image and (c) associated SAED pattern from the thicker TiAlN film. The position of the selected area aperture is marked by a dashed circle in (b). The rings in the SAED pattern have been assigned to hexagonal (solid) and cubic (dashed) nitride phases with larger and shorter lattice spacing, respectively, than the binary compounds. Note that several rings do overlap (dash-dot circles) for both structures.
The Ti2p spectra of TiAlN show four main contributions labelled as A, B, C, and D. This edge originates from 2p to 3d transitions and, hence, shows a L2 (C and D features) and L3 (A and B features) complex structure due to the spin–orbit interaction of the Ti-2p levels. The 3d levels in transition metals can also be split due to crystal-field effects [38]. Typically, the L2,3 splitting in the Ti2p XANES spectrum of c-TiN is 6 eV whereas the splitting due to crystal-field effects is around 1.5 eV [39]. In the case of covalent bonding, such as in TiO2, the L2,3 edges present a well-defined and strong crystal-field splitting (3 eV), with distinct spectral features according to the different coordination geometries [40].
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However, this effect is less pronounced in the case of materials with metallic character [41], such as the case of pure Ti or c-TiN, presenting broader features. Accordingly, the Ti2p XANES spectrum from c-TiN in Fig. 5 shows two main broad features around 459 and 465 eV, in agreement with the results from Soriano et al. [39]. There is also a small contribution of Ti–O bonds from surface oxidation as derived by the peak at 459.5 eV [39]. Soriano et al. [39] has shown that the Ti–O feature is promoted upon progressive surface oxidation until the spectrum resembles that of TiO2. Note that the contribution from Ti–O is absent in the TiAlN films produced with a compound target, related to the higher oxidation resistance of Ti atoms in the ternary nitride as compared to cTiN. Remarkably, the position and relative intensity of the fine structure of the L2,3 edges in TiAlN differ from that of c-TiN. The presence of fine structure suggests a higher contribution from crystal-field effects and points out toward a more covalent bonding configuration around Ti atoms. Interestingly, the relative intensity of features A and B is inversed for [N2] P 25%. The different spectral features for [N2] P 25% should be correlated with the formation of a nanocrystalline hexagonal phase (see Fig. 3). Regarding the Al1s edge, the w-AlN reference sample shows a similar spectrum as that reported in the literature [42,43]. The fine structure in the spectrum is related to transitions from Al-1s to p and d hybridized orbitals, showing a low-energy peak around 1565 eV and a more intense and broader feature around 1571 eV. As predicted theoretically [46,44], the Al1s edge in AlN with rock-salt coordination lacks the low-energy peak characteristic of w-AlN and the more intense feature is downshifted 1 eV with respect to the maxima in the w-AlN spectrum, as indeed observed experimentally [20,21]. Based on this, Al sites with cubic arrangement have been identified in cubic ternary solid solutions [15,17], as highlighted by the reference spectrum for c-Ti1xAlxN (x 0.35) in Fig. 5. In line with the composition and structural analysis, the amorphous and sub-stoichiometric character of the sample grown with [N2] = 10% results in broad features whereas TiAlN samples grown with [N2] P 25% show similar Al1s spectra. In the latter case, the spectral lineshape resembles that of w-AlN as extracted by the presence of a low-energy peak (feature E) followed by a more intense one at higher energies (feature F). Note that also a hump close to 1582 eV is present in the spectra of both TiAlN samples and w-AlN (feature G). The latter feature has also been predicted in ternary wurzite phases and, moreover, its position should not be influenced by the Ti/Al ratio [12] as observed here. In agreement with XRD, the sample grown with [N2] = 25% shows narrower E and F features due to the higher structural ordering (see Table 1). Remarkably, features E and F of TiAlN are shifted to higher energies with respect to w-AlN. This suggests that, although Al sites are in hexagonal arrangements they slightly differ from those at w-AlN. In particular, the peak shifts could result from changes in the wurzite lattice parameters (including bond angle and distance distortions) and/or the bonding character [12] with Ti incorporation. Hence, in conjunction with the changes in the Ti2p spectra, the Al1s edge provides another indirect proof of the formation of a ternary hexagonal phase. The N1s spectra can also yield relevant information about the formation of a ternary compound and deserves special attention, since the N1s edge of ternary cubic and hexagonal nitrides have been previously studied theoretically [12]. On the one hand, the spectrum from c-TiN can be related to transitions from N-1s to unoccupied N-2p states with two clear regions related to the hybridization with the empty Ti-3d (397–404 eV) and Ti-4sp (404–414 eV) bands [45]. The first region shows two components due to the splitting of the Ti-3d states into t2g and eg sub-bands due to crystal-field effects [38]. On the other hand, the lineshape of the N1s spectra in w-AlN corresponds to transitions from N-1s to unoccupied N-2p states [46]. The spectrum shows a clear
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Fig. 5. Ti2p, Al1s, and N1s XANES from TiAlN films grown with different [N2]. The reference spectra from c-TiN and w-AlN compounds grown under similar conditions are displayed too. The Al1s spectrum from a ternary c-Ti1xAlxN solid solution (x = 35) is also included as reference for Al sites in the cubic structure. The spectral intensity has been normalised to one in all cases and the spectra are shifted vertically for clarity.
triple-peak structure whose relative intensity strongly depends on the X-ray incidence angle [47]. The overlap of the N1s features from the reference binary nitrides adds some complexity to the identification between cubic and hexagonal environments in ternary compounds. However, the filling of states around 405 eV (features I–K) in the spectra of TiAlN films indicates the relevant contribution from N sites surrounded by Al atoms in wurzite-like arrangements. The fine structure of the spectrum from w-AlN within 400–410 eV is not present in the TiAlN films. However, it has been predicted [12] that such triple-peak structure is levelled out in wurzite ternary compounds. As in the Al1s edge, the N1s features (I–K) are shifted to higher energies with respect to the positions in w-AlN, which supports the formation of a ternary wurzite phase. There is also a significant contribution from feature H, with energy close to the low-energy peak in the N1s spectrum of c-TiN. This might be interpreted as a certain fraction of N atoms in octahedral coordination. A similar observation was reported by Gîrleanu et al. [48] for Ti0.14Al0.86N films studied by ELNES and attributed to cubic arrangements at grain boundaries. However, in a previous publication [15], the same authors found that a significant fraction of Ti atoms are incorporated in the fourfold sites of the wurzite lattice. Also, the spectrum reported in [48] has different fingerprints as those presented here and, in their case, the triple-peak structure characteristic of w-AlN was clearly resolved. In order to study the N arrangements in more detail, we have first considered the hypothesis that the TiAlN films consist of a mixture of binary phases (that is, excluding the formation of ternary phases). Under such assumption, the N1s spectra of TiAlN in Fig. 5 should be a linear combination of that from c-TiN and wAlN. In order to verify this, the c-TiN spectrum has been subtracted from that of TiAlN and the result is shown in Fig. 6. The subtracted spectrum (solid line) has some similarities with w-AlN (dashed line) but important features are missing, specially the triple-peak structure between 400 and 410 eV. Therefore, we cannot interpret the spectrum of TiAlN films as a sole combination of binary phases. This also agrees with the novel features described for the Ti2p and Al1s edges. Fig. 7 compares our spectra with those predicted for w-Ti1xAlxN phases by Holec et al. [12]. Reproducing the experimental spectra of known structures such as c-TiN and w-AlN is still current
state-of-art research and, obviously, we do not expect to match our experimental spectra where unknown and/or mixed phases may be present. Rather, we are looking for qualitative agreements. Moreover, additional broadening should be applied to the calculations in order to approximate to the experimental conditions [12]. For our comparison, the first peak of the experimental spectrum of w-AlN has been used for energy calibration. First, as a test of the comparison, we can see that there is a good correlation between the measured spectrum of w-AlN and the calculations. Second, the subtracted spectrum from Fig. 6 is found to be similar to that of a w-Ti1xAlxN structure with x = 0.875. This could be interpreted as the result from a mixture between a hexagonal ternary phase with a lower Al content than the nominal one (x = 0.75) and c-TiN. However, it should be noted that no cubic phase has been detected by XRD (Fig. 3) in these samples grown for 1 h (those analysed by XANES). This would imply that the cubic arrangements, if present, are located at the grain boundaries [16]. Taking this into account, the contribution from grain boundaries (and, correspondingly, the c-TiN signal) should increase with [N2] since the crystal-
Fig. 6. N1s spectrum for the TiAlN sample grown with [N2] = 25% after subtraction (solid line) of the c-TiN spectrum and comparison with w-AlN (dashed line).
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in the films. This trend suggests the incorporation of Ti in the hexagonal network, which causes expansion of the lattice. The study of the Ti2p, Al1s, and N1s XANES spectra confirm the formation of a ternary hexagonal compound since the spectra cannot be fitted with a linear combination of the spectra for the binary nitrides and the spectral features are shifted with respect to those of wAlN. Moreover, the previous assumption is supported by comparison with theoretical modelling reported in the literature. Acknowledgments
Fig. 7. Comparison of experimental (dotted curves) N1s spectra of w-AlN and Ti0.25Al0.75N films with calculated edges (thin curves) of w-Ti1xAlxN phases with x = 1 and 0.75, respectively. The subtracted spectrum from Fig. 5 is also compared with a w-Ti1xAlxN phase with x = 0.875. The theoretical curves have been adapted from Ref. [12].
line quality deteriorates (see Fig. 3 and Table 1). However, the N1s spectra in Fig. 5 for TiAlN films grown with [N2] P 25% show a similar relative intensity of feature H (taken as a measure of the c-TiN content in the films). Remarkably, the predicted N1s edge from wTi1xAlxN with x = 0.75 can also be satisfactorily correlated with the fine structure of the experimental N1s XANES spectrum of the Ti0.25Al0.75N film, even including the low energy peak around 398.5 eV (feature H in Fig. 5). Based on the calculations reported in Ref. [12] for Al-rich TiAlN, it would be expected that N sites near Al tend to have a four-coordination (wurzite-like) whereas fivecoordination would be preferred in the vicinity of Ti neighbours. Therefore, we attribute feature H in the N1s spectra of TiAlN films not to cubic arrangements but rather to N atoms surrounded by Ti atoms in hexagonal arrangements. In fact, the N1s spectrum from hexagonal TiN (BN structure) is predicted to have similar features as in the case of c-TiN [12]. It should be mentioned that the calculations in Ref. [12] were limited to the N1s edge and, therefore, it would be desirable to complete the study of the other elemental edges to attain more solid conclusions. Hence, the present experimental results support previous theoretical modelling and could be implemented in further theoretical modelling to understand the formation and electronic structure of ternary nitride phases.
4. Conclusions TiAlN thin films have been grown by reactive DC magnetron sputtering in an Ar/N2 atmosphere from a Ti50Al50 compound target. ERDA-BIC shows that the film composition is Al-rich and that stoichiometric nitrides are produced above a certain threshold of N2 contents in the gas mixture ([N2] P 25%). AFM shows that the surface morphology of the TiAlN films is very smooth (r > 2 nm). In the case of stoichiometric nitrides, glancing-incidence XRD shows that polycrystalline films with wurzite structure are obtained, the c-axis lattice parameter increasing with the Ti/Al ratio
We thank G.S. Fox-Rabinovich for providing cemented carbide inserts and T. Regier for beamline assistance at SGM (CLS). This work is part of the research carried out within the bilateral 2010RU0025 project between CSIC (Spain) and RFBR (Russia). Additional financial support from grants FIS2012-38866-C05-05 (MECO, Spain), AVANSENS no. S2009/PPQ-1642 (CAM, Spain), and CSD2008-00023 (MICINN, Spain) is acknowledged. ERD-BIC measurements have been supported by the EC Integrating Activity ‘‘Support of Public and Industrial Research Using Ion Beam Technology (SPIRIT)’’ under contract no. 227012. The synchrotron work at CLS (Canada) has been supported by NSERC, NRC, CIHR and the University of Saskatchewan. F.S. thanks the EFRE Funds of the European Commission for support of activities within the AME-Lab project. Finally, ARC acknowledges funding from SFRH/BPD/74095/2010 and PTDC/CTM/100756/2008 grants (FTC, Portugal). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
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