Mechanical properties and oxidation resistance of nanocomposite TiN–SiNx physical-vapor-deposited thin films

Surface and Coatings Technology 120–121 (1999) 158–165 www.elsevier.nl/locate/surfcoat

Mechanical properties and oxidation resistance of nanocomposite TiN–SiN physical-vapor-deposited thin films x M. Diserens a,b, J. Patscheider a, *, F. Le´vy b ¨ berlandstr. 129, CH-8600 Du¨bendorf, Switzerland a Eidgeno¨ssische Materialpru¨fungs- und Forschungsanstalt (EMPA), U b Institut de Physique Applique´e, EPFL, CH-1015 Lausanne, Switzerland

Abstract Thin films of TiN–SiN have been prepared by reactive unbalanced magnetron sputtering from two opposite Ti and Si targets. x The silicon concentration in the deposited coatings is varied between 0 and 16 at.%. The deposited films are composed of TiN nanocrystallites embedded in an amorphous SiN matrix. These nanocomposite coatings exhibit improved mechanical properties x in comparison with TiN deposited under the same conditions. Whereas the hardness measured by nanoindentation is about 27 GPa for TiN, it reaches 38 GPa in TiN–SiN films containing 5 at.% Si and decreases to the values of amorphous SiN at x x silicon concentrations above 15 at.%. Besides higher hardness values and improved wear resistance, these composite coatings are superior to TiN in their resistance against oxidation. The oxidation resistance is gradually enhanced by increasing the silicon concentration in the films. At 800°C in air, TiN–SiN films with an Si content as low as 5 at.% exhibit a one order of magnitude x lower oxidation rate compared with that of TiN. This oxidation resistance improvement is explained by the presence of the amorphous SiN phase at the TiN grain boundaries, which limits the oxidation in these high diffusion paths and prevents x recrystallization of TiO . © 1999 Elsevier Science S.A. All rights reserved. 2

1. Introduction Titanium nitride coatings are widely used to improve the wear resistance and durability of cutting tools, drills or mills because of their high hardness (>20 GPa) [1]. However, in addition to good mechanical and adhesion properties, good chemical stability is also required for protective coatings. Since TiN starts to oxidize severely at temperatures as low as 500°C, efforts have been undertaken to improve this shortcoming. The addition of aluminum to TiN improves the oxidation resistance up to 800°C and an increased hardness is observed [2,3]. The enhanced oxidation resistance is a consequence of the formation of an aluminum-rich protective passive layer at the surface [4–6 ]. Like aluminum, silicon is a light element that forms a stable oxide and for this reason its addition to TiN has been investigated. In accordance with the Ti–Si–N phase diagram, which does not present any stable ternary phase under equilibrium conditions [7], two-phase TiN–Si N coatings were produced by chemical vapor 3 4 deposition (CVD) at high temperature (above 1050°C ) * Corresponding author. E-mail address: [email protected] (J. Patscheider)

by Hirai et al. [8] in 1982. Later, Li et al. [9] and Veprˇek and coworkers [10–12] deposited Ti–Si–N films by plasma-assisted CVD (PACVD), which led to a reduction of the deposition temperature to 550°C. These films were identified as nanocomposites consisting of nanometer-sized TiN crystallites surrounded by an amorphous Si N matrix. The TiN–Si N coatings exhibited promis3 4 3 4 ing mechanical properties (hardness of 50 GPa and elastic modulus of 500 GPa at an Si content of 8 at.%) [9–12] and good wear resistance [13]. Deposition of these nanocomposite coatings by PACVD using TiCl 4 and SiCl or SiH are commercially problematic because 4 4 of the safety hazards imposed by the volatile chemicals used. The use of chlorine-containing species is an additional drawback due to the corrosive action of unreacted chlorine incorporated in the coatings [14,15]. These two shortcomings were our incentives to deposit TiN–Si N nanocomposite coatings by a physical vapor 3 4 deposition (PVD) process, which avoids the hazardous gases utilized in the PACVD process. The use of reactive unbalanced magnetron sputtering has been demonstrated to be a successful technique to deposit such coatings [16 ]. Similar results have been obtained using r.f. magnetron sputtering from separate targets [17,18]. For electronic devices, thin films of Ti–Si–N have been

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produced as diffusion barrier layers for Cu or Al metallization by reactive co-sputtering from Ti–Si composite targets [19,20]. These silicon-rich films are electrically conductive, show good diffusion barrier properties and have either an amorphous or a nanocrystalline structure. The aim of the present work is to widen the understanding of this new material and to examine two of its most important properties for applications as protective coatings, namely the hardness and the oxidation resistance.

2. Experimental procedure 2.1. Films deposition The TiN–SiN thin films are deposited by reactive x unbalanced magnetron sputtering. The deposition system (Fig. 1) is a cylindrical PVD reactor with two opposed titanium and silicon targets. The use of these separate targets allows an independent regulation of each source and enables the adjustment of the Ti:Si ratio in the gas phase. The unbalanced magnetrons are arranged in a closed-field configuration to provide both a high sputtering rate on the targets and a high plasma density in the zone of deposition [21,22]. The planetary rotation of the sample holder provides a uniform exposure of the growing films to the two targets. An r.f. (13.56 MHz) bias voltage between −100 V and −125 V was applied to the samples during the deposition. The

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base pressure in the reactor was lower than 2×10−4 Pa. During the deposition, the total pressure was measured with a capacitance manometer. All the experiments were carried out with 25 sccm argon flow leading to a partial pressure of 0.5 Pa. The power densities at the titanium target (4 W/cm2) and at the silicon target (1.2 W/cm2) were kept constant for all depositions. The variation of the silicon content was achieved by using different nitrogen partial pressures, which were varied between 8×10−2 and 1.2×10−1 Pa, and by variation of the substrate bias voltages. The silicon concentration in the films was determined by X-ray photoelectron spectroscopy ( XPS) and referenced to by Rutherford backscattering standards; it varied between 1 and 16 at.%. Under magnetron discharge conditions, the reactivity of silicon with nitrogen is substantially lower than that of titanium. For this reason, the nitrogen flow must be introduced at the silicon target to achieve a good nitridation of both elements. In this way, the nitrogen content obtained varied between 43 and 51 at.%. At a deposition rate of about 0.5 mm/h, 2 mm thick layers were deposited on hardened 1.2436 steel substrates with a 300 nm Ti interlayer to guarantee the adhesion of the coatings during the oxidation and the mechanical test experiments. Because of the planetary rotation of the substrates the measurement of sample temperature is difficult. The temperature was instead measured in the center of the chamber with a steel-encapsulated thermocouple; it never exceeded 350°C. Calibration measurements with temperature-indicating paints showed that the readout of the thermocouple, even at bias values above 200 V, was accurate within 20°C. X-ray diffraction ( XRD), XPS, transmission and scanning electron microscopy ( TEM and SEM ) investigations have shown that the deposited coatings are nanocomposites consisting of TiN nanocrystallites (grain size smaller than 10 nm at an Si content exceeding 4 at.%) in an amorphous SiN matrix [16 ]. x 2.2. Nanoindentation procedure

Fig. 1. Closed-field unbalanced magnetron sputtering system with two opposed DC-powered targets and a rotating sample holder with planetary motion.

All indentation experiments were performed using a Nano Indenter XP with a Berkovich tip (a three-sided pyramid with the same area-to-depth ratio as a Vickers indenter). By superimposing an alternating force to the conventional indentation loading ramp, this setup allowed the measurement of the contact stiffness at each depth [23]. This continuous stiffness measurement permits the determination of a depth profile of the nanohardness H and of the Young’s modulus E throughout the coating within only one indentation cycle. In order to obtain mean and error values, nine indentations were performed on each coating. The calibration of the system was obtained by optimizing the area function of the diamond tip in the indentation depth range between 20

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Fig. 2. Determination of nanohardness and Young’s modulus by continuous stiffness measurement: example on a TiN–SiN coating containing x 5 at.% silicon.

and 500 nm on a fused silica reference sample (H= 9.3 GPa and E=72 GPa). The intrinsic mechanical properties of these coatings were measured at indentation depths between 100 and 200 nm ( Fig. 2). On the one hand, in this range the contact stiffness measurement is not affected by the surface roughness, as is the case in the first tenths of nanometers; on the other hand, the deformation zone induced by the indentation does not reach the film– substrate interface. The influence of the mechanical properties of the steel substrate (H=8 GPa and E= 200 GPa) on the measurements is therefore avoided. 2.3. Oxidation resistance measurements To examine the oxidation resistance, the coated samples were exposed to a heat treatment at fixed temperatures between 700 and 1000°C in air. After a 10 min heating ramp, the temperature was kept constant for 2 h and cooled down within 10 min. Oxide thicknesses greater than 200 nm were determined by using a CSEM Calo-Wear ball cratering system with a 25 mm steel ball and an abrasive suspension of 0.25 mm diamond particles in alcohol. On the basis of microscopic measurements of the wear crater, the oxide thickness is calculated. To determine lower oxide thicknesses, Auger electron spectroscopy depth profiling is performed with a 3 keV Ar+ ion beam. The sputtering rate was calibrated on a 300 nm thick oxide layer grown on a TiN–SiN coating (Si content: 14 at.%). x 3. Results and discussion 3.1. Mechanical properties The nanoindentation measurements show that even a low level of silicon incorporation in TiN coatings

modifies its mechanical properties (Figs. 3 and 4). TiN–SiN with an Si content as low as 1 at.% already x exhibits a strong hardness increase in comparison with TiN deposited under the same conditions. At about 5 at.% Si, the hardness of TiN–SiN reaches a maximum x of 38 GPa. The hardness increase is due to the strong grain refinement [11], which implies a high grain boundary density. It is known from PACVD coatings that the maximum hardness coincides with a minimal grain size, which implies a high density of grain boundaries [10]. These grain boundaries hinder the dislocation migration in the TiN crystalline phase and hence limit the plastic deformation induced by the indentation. Crack propagation in the amorphous phase is also blocked by the presence of TiN crystallites [24]. Both hardness and elastic modulus decrease in TiN–SiN above 6 at.% silicon. At these compositions, x the TiN grain size refinement below 5 nm leads to mechanical properties that are comparable to those of amorphous silicon nitride in spite of the fact that the structure stays composite. One probable explanation is that the TiN grain size is small enough to permit crack propagation around the crystallites and, therefore, the blocking effect of grain boundaries is no longer effective. The hardness measurements (Fig. 4) also illustrate that a complete nitridation is necessary to achieve good mechanical properties. From the lattice parameters measured by XRD and the Ti 2p binding energies mea3/2 sured by XPS, we conclude that titanium is fully nitrided, a known prerequisite for optimal hardness of TiN [1]. This also suggests that the silicon nitridation is of great importance and for this reason the argon introduction at the silicon target is suppressed in order to increase the nitrogen partial pressure gradient between the two targets. In this way, the reactivity differences between titanium and silicon are taken into account and a low silicon nitridation in the films is avoided.

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Fig. 3. Young’s modulus of TiN–SiN coatings as determined by nanoindentation versus silicon content. x

Fig. 4. Nanohardness of TiN–SiN coatings as function of silicon and nitrogen contents. x

3.2. Oxidation resistance The oxidation of TiN is limited by the diffusion of oxygen [25,26 ] and therefore the oxide thickness d oxide can be expressed by a parabolic law: d =앀2Dt, oxide where t is the oxidation time and D the diffusion coefficient of oxygen in TiO . As the oxidation of TiN 2 is a thermally activated process, D can be described by the Arrhenius law: D=D e−Ea/RT, 0 where E is the thermal activation energy associated a

with the parabolic oxidation process and D the pre0 exponential factor. The activation energy for the oxidation of TiN in the temperature range between 500 and 1200°C reported by several authors varies from 187 to 198 kJ/mol (1.94–2.05 eV/molecule) [25–27]. The differences can be related to differences in morphology, defect density and stoichiometry. In this work, we performed only one TiN oxidation test at 700°C as a reference. The measured oxide thicknesses are in the expected range [25–27] (Fig. 5). The silicon incorporation in TiN improves the oxidation resistance significantly. The oxidation rate of the TiN–SiN coatings with highest hardness (Si content: x 5 at.%) is lowered by approximately an order of magni-

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Fig. 5. Arrhenius plot of the oxide layer thickness after a 2 h heat treatment. The symbols indicate our measurements on TiN and TiN–SiN x coatings and are compared with data from the literature [4,25,26 ].

tude at 800°C compared with TiN. The apparent activation energy for the oxidation of this film is 268 kJ/mol (2.78 eV/molecule). At 14 at.% silicon content, TiN–SiN has an even better oxidation resistance than x Ti Al N [4]. The oxidation experiments at different 0.5 0.5 temperatures on this silicon-rich coating suggest that the oxidation of TiN–SiN proceeds by two different x mechanisms in the temperature regime investigated. Below approximately 820°C the oxidation proceeds very slowly. This indicates that TiN–SiN forms a stable x passive layer at low temperatures limiting oxygen diffusion. An oxide layer of maximum thickness should be reached after a prolonged oxidation time, as is observed on TiN–Ti Si coatings [27]. Above 820°C, the second 3 5 oxidation mode of this silicon-rich TiN–SiN follows an x Arrhenius law with an apparent activation energy of

293 kJ/mol (3.03 eV/molecule) ( Fig. 5). The XPS depth profile throughout the oxide layer formed on a 10 at.% Si coating at 700°C ( Fig. 6) shows the presence of SiO and of TiO at the surface; no silicon enrichment 2 2 in the oxide layer is observed. This is a marked difference to the oxidation behavior of Ti Al N, where the 0.5 0.5 passive layer formation is associated with a strong aluminum migration to the surface [4–6,28]. Atomic force microscope (AFM ) images of the surface topography of the oxide scales formed on TiN–SiN containing 11 at.% Si show the changes in x surface morphology and roughness (Fig. 7). By increasing the oxidation temperature it is observed that the TiN–SiN finely structured surface first becomes x smoother with the formation of a passive layer at 770°C. Increasing the temperature above 850°C leads to a

Fig. 6. XPS depth profile through the oxide layer formed after 2 h at 700°C on a 10 at.% Si TiN–SiN coating. x

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Fig. 7. 2mm×2mm non-contact AFM images of oxidized TiN–SiN coatings (14 at.% Si) after a 2 h oxidation stage in air at different temperatures: x (a) no heat treatment; (b) 770°C; (c) 860°C; (d ) 880°C; (e) 940°C. The mean roughness R measured on a 100 mm2 area equals: (a) 7 nm; (b) 5 nm; a (c) 5 nm; (d ) 30 nm; (e) 100 nm.

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significant increase in surface roughness, which finally leads to the formation of facetted grains at 940°C. Grazing-angle XRD indicates the presence of wellcrystallized TiO in the rutile form, whereas no signals 2 from crystalline SiO phases are observed ( Fig. 8). 2 Combining the kinetic data from Fig. 5 with the clear evidence of SiO at the surface from XPS and the 2 topographical information from the AFM observations allows us to propose an oxidation mechanism for TiN–SiN coatings that takes into account the composite x nature of these coatings: below about 850°C amorphous silicon nitride, which separates the crystalline TiN grains, acts as an efficient diffusion barrier against oxygen diffusion and in this way protects the TiN grains from oxidation. For comparison, silicon nitride is used in microelectronic devices as a diffusion barrier [29]; this behavior as a diffusion barrier is supported by the concentration-dependent oxidation rate where samples with high silicon nitride contents showed the slowest oxidation. This diffusion-controlled mechanism has a smaller apparent activation energy than that observed for the oxidation of TiN. The apparent activation energy of the second stage of oxidation above 850°C amounts to about the same value as observed in the oxidation of silicon-free TiN [25–27]. This regime of increased oxidation rate is characterized by the formation of large grains of rutile. It is meaningful to assume that the diffusion of Ti atoms through the amorphous silicon nitride and/or silicon oxide layer(s) is the rate-limiting step for the formation of rutile, once the amorphous diffusion barrier is no longer effective. The observed activation energy apparently characterizes the recrystallization of rutile forming from TiN and oxygen without

diffusion barriers as proposed by Wittmer et al. [25] and Lefort et al. [26 ]. The compressive internal stresses (3 GPa) measured by the sin2(Y ) XRD method [30,31] of the as-deposited films are released during the heat treatment. This stress relaxation permits both the precise determination of the lattice parameters of TiN located in the TiN–SiN x coating and of TiO (rutile) in the oxide layer. The 2 ˚ with the values obtained ( Fig. 8) agree within 0.005 A tabulated values for TiN and rutile [32]. This agreement virtually excludes the possibility of silicon incorporation as a substitutional or interstitial element in both phases (for comparison, aluminum has approximately the same atomic radius as silicon and its substitution in the facecentered cubic structure of TiN induces a substantial lattice parameter reduction — a( Ti)=0.424 nm and a( Ti Al N )=0.416 nm [33] — in spite of the absence 0.5 0.5 of silicon-containing diffracting phases. Accordingly, the oxide layer consists of a crystalline-rutile–amorphoussilica composite, in agreement with the Ti–Si–O system phase diagram at 1000°C, which does not exhibit any oxygen-rich ternary phase [34]. The amorphous silica is known for its anti-diffusion properties and its presence at the TiO grain boundaries 2 maintains the stability of the TiN–SiN passive layer at x low oxidation temperature by hindering oxygen and titanium atom migration. Above a threshold temperature, depending on the silicon content, the barrier effect of amorphous silica can no longer prevent Ti diffusion. Consequently, rutile recrystallization occurs, which induces a strong roughening of the surface and acceleration of the oxidation process.

Fig. 8. Grazing incidence X-ray diffractogram of a TiN–SiN coatings (14 at.% Si) after 2 h oxidation at 943°C showing the presence of TiN and x TiO (rutile) crystallites. 2

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4. Conclusion Nanocomposite TiN–SiN films have been deposited x by reactive unbalanced magnetron sputtering. Codeposition of silicon leads to a strong refinement of the TiN grain size in the nanometer range while silicon nitride forms an amorphous phase. The small grain size of TiN induces a high density of grain boundaries between the crystallites and the amorphous SiN . Under x stress, these interfaces act as barriers against dislocation migration and, in this way, increase the barrier for plastic deformation. For these reasons, TiN–SiN coatx ings exhibit a high hardness with a maximum of about 38 GPa at a silicon content of 5 at.%. The oxidation resistance of TiN is gradually enhanced by silicon incorporation. Even a level of 5 at.% leads to a ten times lower oxidation rate compared with siliconfree TiN. Kinetic experiments indicate an oxidation mechanism involving two steps. In the temperature regime below 850°C the oxidation proceeds slowly because the amorphous silicon nitride phase reduces the oxygen diffusion at the grain boundaries. Above this temperature the oxidation of TiN–SiN is dominated by x recrystallization of TiO (rutile) in the stable oxide layer. 2 Both the enhanced hardness and the high oxidation stability of composite TiN–SiN coatings are a consex quence of the massive presence of grain boundaries and hence of the nanometer-scaled microstructure of the TiN–SiN coatings. x

Acknowledgement The support of this work by contributions of both the presidents’ funds of EMPA and EPFL is gratefully acknowledged.

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