Oxidation Behavior and Mechanisms of TiAlN/VN Coatings
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- Zhou, Z., Rainforth, W., Rodenburg, C. et al. Metall and Mat Trans A (2007) 38: 2464. doi:10.1007/s11661-007-9293-4
Hard wear-resistant coatings require excellent oxidation resistance for high-speed machining operations. Moreover, the oxide formed is integral to the frictional behavior and therefore the success of the coating. The oxidation behavior of TiAlN/VN nanoscale multilayer coatings was investigated using high-resolution techniques and was compared with TiN and TiAlN coatings. Static oxidation of TiAlN/VN films was studied in the range 550 °C to 700 °C, and characterized by high-temperature in-situ X-ray diffraction (XRD) and scanning transmission electron microscopy/energy-dispersive X-ray/electron energy loss spectroscopy (STEM/EDX/EELS) of selected surface cross sections. The oxidation resistance of TiAlN/VN was found to be controlled by the VN layers, and consequently, oxidation was initiated at a lower temperature than TiN and TiAlN coatings. The onset of oxidation of the TiAlN/VN coating was found to be ≥550 °C with the VN being the first component to oxidize. At temperatures >600 °C, a duplex oxide structure was formed; the inner layer comprised a porous region of Ti-rich and V-rich nanocrystallites, while several phases were observed in the outer region, including V2O5, TiO2, and AlVO4. V2O5 was the dominant oxide at the outer layer at ≥638 °C. The outward diffusion of V depended on the species present; in the inner layer, V was present as V3+, V4+, whereas a significant V5+ was dominant in the outer layer of oxide at ≥638 °C. An Au marker study suggested roughly equal diffusivity of cations outward, and oxygen inward diffusion occurred during oxidation.
Titanium nitride (TiN) with the B1 NaCl structure has been widely used as a hard wear-protective coating since the 1980s. However, a major limitation of TiN coatings for high-speed machining applications is that they oxidize rapidly to rutile, TiO2, at temperatures above 500 °C. Initial oxidation is parabolic, with O diffusion through TiO2 to the oxide/nitride interface as the rate-limiting step. The large difference in molar volumes of TiO2 and TiN results in spallation of the oxide scale, resulting in a transition from parabolic to pseudo-linear kinetics.[2,3]
TiAlN coatings provide much improved high-temperature oxidation resistance, up to 750 °C to 900 °C, and have consequently been used extensively for high-temperature cutting operations with minimum use of lubricant or dry machining. Titanium aluminum nitride (Ti1-xAlxN) has the B1 NaCl structure with 0 ≤ x ≤ 0.52, but as x is increased, phase decomposition occurs into B1 NaCl and wurtzite AlN structures.[5,6] Ti0.5Al0.5N exhibits stable parabolic oxidation kinetics, associated with an oxide comprised of a passive double layer oxide, which is a result of outward diffusion of Al to form Al-rich oxide at the topmost surface and inward diffusion of O to form Ti-rich oxide at the interface to TiAlN. However, during sliding wear tests, the friction coefficient of TiAlN is usually high, typically μ = 0.8 for TiAlN measured from the pin-on-disc test against alumina ball.
Both the friction and wear performance of TiAlN can be improved through the quaternary addition of V, as monolithic Ti-Al-V-N coatings or as TiAlN/VN multilayers. The VN, which also has the B1 NaCl structure, exhibits continuous solid solubility with a number of metal nitrides and carbides, such as TiN, TiC, TiAlN, VC, and NbN. Knotek et al. found that the addition of V to TiN and Ti-Al-N enhanced wear performance, with the best wear resistance obtained in a Ti-V-N coating with 29 at. pct V. TiAlN/VN multilayers, which incorporate VN into TiAlN as a nanoscale multilayer (also referred to as a “superlattice”) coating, have exhibited excellent sliding wear resistance (1.26 × 10−17 m3·N−1 m−1) with a lower friction coefficient at room temperature (μ = 0.4, pin-on-disc test, Al2O3 ball counterpart) in comparison to other wear-resistant coatings, e.g., TiAlN/CrN (sliding wear coefficient 2.4 × 10−16 m3·N−1 m−1, μ = 0.7 to 0.9).
The oxidation of VN has been investigated by several authors.[11–13] Oxidation is complex, with columnar V2O5 being the dominant scale, but also suboxides V6O13, VO2, and the oxygen-deficient Magnéli phases, VnO2n−1, VnO3n−1, or VnO3n−2, which have the rutile structure (TinO2n−1) formed at the interface with a lamellar structure. The limiting step was found to be the diffusion of both anionic and cationic species through the V2O5. Not surprisingly, the reaction rate accelerated dramatically at the melting point of the V2O5 (674 °C), where oxidation became catastrophic.
Frictional heating during sliding often results in the formation of surface oxides. These oxides can have either a beneficial or detrimental effect on wear and friction behavior.[14,15] In the case of TiAlN/VN, it is believed that the products of oxidation can substantially reduce friction. For example, high-temperature wear tests of TiAlN/VN coatings and Ti-Al-V-N films with various contents of V (2 to 25 at. pct in targets) against an Al2O3 ball showed a dramatic decrease of friction coefficient from >0.8 at 500 °C to 0.2 at 700 °C. Gassner et al. provided some evidence that this is associated with the formation of a variety of Magnéli phases, while V2O5 has also been observed in room-temperature tests using Raman spectroscopy. The V2O5 has a melting point of 674 °C and inherently low friction, leading to suggestions that the formation of this oxide is the key to understanding the low frictional behavior of these coatings. Moreover, the Magnéli phases, VnO2n−1, VnO3n−1, or VnO3n−2, all contain crystallographic shear (CS) planes, which are known to promote low friction.
It is clear that the oxidation behavior, both under static and sliding conditions, strongly affects the performance of these coatings. To date, very little knowledge is available on the static oxidation of the Ti-Al-V-N system. Preliminary investigations by the current authors into the thermal behavior of TiAlN/VN coatings[22,23] showed that oxidation initiated at ≥550 °C with V2O5, TiO2, and AlVO4 identified at 600 °C. Small changes in (Ti + Al)/V ratio changed the exact temperature for the onset of oxidation, with an increase in (Ti + Al)/V increasing the onset temperature, but these changes did not apparently change the mechanism.[23–25] The current work provides a comprehensive understanding of the oxidation mechanisms in these coatings for approximately constant (Ti + Al)/V. Although not the focus of the study, oxidation kinetic data are included for TiN and TiAlN coatings as well, to provide a baseline.
Details of Coating Deposition Conditions, Composition, Dominant Crystallographic Texture, and Residual Stress
Al (At. Pct)
Ti (At. Pct)
V (At. Pct)
(Ti + Al)/V
Residual Stress (GPa)
In order to compare the oxidation kinetics of the TiAlN/VN multilayer coatings, TiN and TiAlN coatings were prepared. The coatings were deposited on the same machine using bias voltages of 75 V. Full details of the deposition and structure of these coatings are provided elsewhere.[27,28]
A Cahn TG 131 microbalance was used for the thermogravimetry work, with the following specifications: temperature drift stability of 10 μg/°C, temperature repeatability of ±3 °C, and mass sensitivity of 1 μg. The TG allows the determination of the onset point of oxidation and the extent of weight gain. All coatings were deposited on stainless steel 304 coupons (dimensions 50 × 15 × 0.75 mm). The coupons were completely covered with coating material, including the hole drilled for fixing the sample in the TG furnace, such that the oxidation would only occur in the coating and the substrate would only become oxidized when the coating failed. Oxidation of TiAlN/VN was assessed dynamically by thermogravimetric analysis (TGA) in a linear-temperature-ramp (400 °C to 1000 °C at 1 °C/min) mode. Isothermal heat treatment at 550 °C and 600 °C was also conducted. Coatings were subsequently heat treated at 550 °C, 600 °C, 638 °C, and 672 °C in a conventional ambient air atmosphere furnace in order to correlate the TG traces with phase transformations on the coating surfaces, in particular, the formation of V2O5. These temperatures were selected based on the TG trace.
The X-ray diffraction (XRD) of the as-deposited and oxidized coatings was undertaken on a PHILIPS MRD-Xpert, Cu Kα radiation with 2θ from 10 to 140 deg using standard Bragg Brentano diffraction conditions. Glancing angle geometry with 1 deg incidence was used when the oxide scale was thin, i.e., at 550 °C.
A coating deposited with a bias voltage of −85 V was examined using high-temperature XRD. The sample was heated in a Bruker D5005 diffractometer with a high-temperature stage from room temperature to 700 °C at 25 °C/min. The sample was held at the chosen temperature for 40 minutes to complete a scan. The θ to 2θ scan was performed from 10 to 55 deg, at 0.014-deg step size using Cu Kα radiation.
Experimental Conditions for the Acquisition of Elemental and Jump Ratio Images on the Gatan GIF200
N-K (401 eV)
Ti-L2,3 (456 eV)
V-L2,3 (513 eV)
O-K (532 eV)
N/A (jump ratio)
Slit width (eV)
The Au marker studies were undertaken to provide information of the relative diffusion rates of the cations and anions. A TiAlN/VN coating was covered by a piece of net fabric and then sputter coated with Au for 12 minutes. The Au islands were left on the coating surface once the net fabric was removed. The coating was then heat treated in air at 600 °C for 0.5 hours. It was assumed that Au does not react with the coating and the oxides up to this temperature. The oxidized coating was then Ni plated to preserve oxide scale prior to cross-sectional metallographic preparation.
3.1 Composition and Structure of the As-Deposited Coatings
The coatings deposited are shown in Table I. Relative contents of (Ti + Al)/V were determined by SEM/EDX at the plan-view coating surface, using a standardless routine (error ±0.5 pct), and therefore the results should be regarded as semiquantitative. Coatings deposited at −75 and −85 V had approximately the same (Ti + Al)/V ratio. Low-angle XRD revealed multilayer repeat periods of ∼3 nm for all coatings. The coating deposited with a substrate bias voltage of −85 V possessed a residual compressive stress of −7.9 GPa and (111) texture, whereas the coatings deposited at −75 V bias had a lower residual stress of −3.9 GPa and exhibited a (110) texture. More details about the influence of substrate bias voltage on the coating’s composition, microstructure, texture, residual stress, and hardness were provided in Reference 30.
3.2 Continuous and Isothermal TGA
Figure 2(b) gives the isothermal TG traces obtained by heating the TiAlN/VN coating (−85 V bias) for 1.5 hours at 550 °C and 600 °C, respectively. The temperature profile is also shown, from which it is clear that oxidation may have started during the heating phase for the 600 °C sample. The oxidation rates at 550 °C initially appeared logarithmic, but changed to parabolic after ∼0.4 hours, and became linear after ∼0.7 hours with a much slower rate of mass gain. The behavior of the 600 °C sample was similar, except the parabolic oxidation rate was more rapid, and the oxidation rate >0.7 hours was slightly greater (0.08 g/m2 h) compared to the 550 °C sample (0.06 g/m2 h). The oxidation kinetics of the TiAlN/VN coatings at 550 °C and 600 °C closely resembles that of the Nb and Ta metals below 600 °C.
3.3 Surface Structure and Phase Evolution during Oxidation
3.4 Cross Sections of Oxidized Surface by TEM
The first notable point is that the N content in the oxide film appears to be very small, suggesting that most of the N must have been liberated as gas, discussed later. However, the EEL spectra show that the structure of the N K edge is substantially different at the interface compared to the coating. The shape of the edge (near edge structure) reflects the bonding environment of the N and indicates a significant change in bonding at the oxide interface.
The energy-filtered maps indicate an abrupt interface between the oxide and the coating. The O K edge is distorted by the overlying near edge structure from the V L2,3 edge. However, it would appear that while O is clearly present at the interface, it is below detectable quantities 3 nm into the coating. The V L2,3 intensity significantly reduced across the interface and was generally lower within this oxide layer, although the exact intensity changed from point to point depended on the local phase constitution, as shown by the V energy filtered map in Figures 13 and 14. The reduction in V intensity was coincident with the higher porosity in the coating, but this did not reduce the intensity in the Ti L2,3 or Al L2,3 maps. The intensity in the Ti L2,3 maps also varied substantially from point to point, although locally in the oxide film the edge intensity was as high as in the coating, as in the spectrum shown in Figure 15.
Thus, the TEM section indicates that at 638 °C, the oxide comprises three layers: two layers with essentially the same phase constitution as at 600 °C, but with a change in proportion of phases, and with the V2O5 forming a near continuous outer layer, whereas this layer was occasional at 600 °C.
3.5 Au Marker Study
4.1 Oxidation Kinetics and the Temperature for the Onset of Oxidation
There was no evidence that the oxidation kinetics or mechanism of the TiAlN/VN was significantly modified by the substrate bias voltage during deposition, even though the bias voltage did appreciably change the crystallographic texture and residual stresses (Table I). Thus, the observations of oxidation mechanism from the bias voltages of −75 and −85 V can be combined.
The TGA traces of TiN and TiAlN coatings were believed to be, to a good approximation, representative of the mass gain of the coatings, because there was no significant diffusion between TiAlN coatings and stainless steel substrate was expected up to 600 °C. However, minor Fe has been found at columnar grain boundaries of TiAlCrN at 700 °C resulting from the diffusion of Fe from the substrate. In the case with TiAlN/VN, the mass gain in TGA probably included a contribution from the steel substrate at 600 °C, because V2O5 forms eutectics FeVO4 with Fe2O3 at 635 °C and 840 °C corresponding to phase diagram V2O5/Fe2O3. However, the TEM examination presented here indicated no significant ingress of Fe into the coating at the temperatures studied here. The other onset of rapid oxidation at 840 °C was possibly associated with oxidation of substrate, but at that temperature, the coating had completely failed and was therefore of no interest.
The onset of rapid oxidation of the TiAlN/VN coating was around 635 °C (Figure 2(a)), which compares to ∼700 °C for TiN and ∼900 °C for TiAlN. Thus, as expected, the addition of VN to TiAlN reduced the oxidation resistance. However, as noted earlier, the product oxides are known to be beneficial during sliding contact, and therefore, this should not necessarily be regarded as a disadvantage for tribological applications. Although no data were obtained for VN in the current study, Borgianni et al. have shown that VN oxidizes at just below 400 °C in flowing oxygen, clearly indicating that the oxidation kinetics of TiAlN/VN was dictated by a combination of the TiAlN and VN, not just the VN.
The oxidation of the TiAlN/VN appeared to have parabolic stages at both 550 °C and 600 °C (Figure 2(b)), although the mass gain after 0.7 hours decreased abruptly to a much lower rate. Note that this transition was not believed to be associated with exposure of the substrate, because TEM sections through the coating indicated that oxidation was wholly contained within the coating well after this change in kinetics. The kinetics could not be adequately evaluated at 638 °C and 670 °C because of rapid failure of the coating through oxidation. The oxidation kinetics of TiAlN have been found to be parabolic (e.g., Reference 7) also with a change in oxidation kinetics to lower growth rates after a certain incubation time, which was believed to be due to the formation of a passive aluminum oxide layer. In contrast, the isothermal oxidation of VN at 640 °C has been found to follow linear kinetics, rather than the parabolic kinetics observed here, indicating that although the addition of VN to TiAlN had significantly adjusted the oxidation kinetics, it had not prevented the formation of a passivating oxide layer.
4.2 Oxide/Coating Interface
The interface between the oxide and the substrate was abrupt (Figures 12 and 17). The oxide front did not follow the multilayer structure, as shown by the orientation of the Fresnel fringes (which delineate the position of the individual layers), which could be imaged right up to the oxide/coating interface.
The oxide adjacent to the coating contained extensive porosity, with a near continuous layer at the interface itself (confirmed by the presence of Fresnel fringes around pores and by the abrupt changes in contrast in the thickness map in Figure 13). The rounded shape of the pores suggests they did not originate from microcracks arising from stresses induced by changes in the molecular volume on oxidation. The porosity is consistent with the reaction of the nitride coating with oxygen to liberate nitrogen gas. Point EEL spectra were obtained from this region (Figure 15). The shape of the N K edge was appreciably different to that found in the coating, with a single peak found at the interface, compared to a double peak in the coating. Moreover, the position of the first peak for the interface K edge was shifted by ∼1.4 eV with respect to the first peak for the K edge for the spectrum from the coating. The interface N K edge can be interpreted in two ways. First, the peak shape and first peak intensity fits that of V2N, as reported by Hofer et al. However, they observed a shift of ∼1 eV between the first peak for VN and V2N, compared to the ∼1.4 eV observed in the present work. Alternatively, the N K edge could be interpreted as nitrogen gas, as shown by Trasobares et al., who examined nitrogen gas inside a carbon nanotube and observed a single peak N K edge with the peak located at 401 eV. In the present work, the peak shape and the increased intensity at the interface in the V L2,3 map (Figure 13) are more consistent with V2N, while the peak position is more appropriate to nitrogen gas. The porosity at the interface was most probably a result of the liberation of nitrogen gas by oxidation of the coating (as discussed later). Moreover, the porosity was typically 5 to 20 nm in diameter and could have been contained within the TEM sample, which was probably of the order of 30-nm thickness in this region. However, while the results do not uniquely identify the presence of nitrogen gas or V2N, it is clear that the bonding of the nitrogen in this region is completely different than that in the coating.
The Fresnel contrast present in the coating right up to the interface (Figures 12 and 17) demonstrated that there was no preferential oxidation of VN in preference to the TiAlN. In this respect, the TiAlN did not provide any “protection” to the VN layers. Moreover, although the TiAlN was present as a discrete phase, the onset of oxidation of this phase had been reduced by >100 °C from the ∼750 °C observed in monolithic TiAl(Cr)N coatings (Figure 2). This is probably because of the elemental intermixing between TiAl and V in the multilayers, such that all TiAlN layers contained appreciable quantities of V within them. Equally, the onset of oxidation of the TiAlN/VN occurred at a much higher temperature than for VN (∼400 °C in flowing oxygen), suggesting that the Ti and Al content of the VN layers was also significant in raising the oxidation resistance.
As shown in Figures 12 and 17, the oxide grew preferentially along columnar grain boundaries. However, the EFTEM maps and associated point EELS analyses showed that the oxygen had not penetrated far into the coating. The structure of columnar grain boundaries in PVD coatings is a strong function of the deposition conditions. Boundaries can exhibit voiding, the volume fraction of which increases with increasing heat-treatment temperature. However, in the current study, most of the columnar boundaries were sufficiently dense and the composition homogeneous, such that they did not appreciably affect the oxidation kinetics.
4.3 Reactions/Stabilities of Phases
The increase in oxygen potential as the oxide grew would be expected to transform any VO or VO2 in the series VO → VO2 → V2O5.
Gold–Schmidt Ionic Radius (nm)
ΔG1000, kJ, 727 °C
Melting point, °C
Molecular volume, cm3
18.8 (TiN: 11.8)
The Au marker studies were used to give some indication of the relative diffusion rates of the cations and anions, given the absence of reliable diffusion data. A comparison of the oxide with (Figure 18) and without the Au marker indicated that the marker had not altered the morphology and thickness of the oxide. The marker for the 600 °C, 0.5-hour sample was located at the interface between the inner and oxide scales. This suggests roughly equal fluxes of the outward diffusion of the cations to form the outer layer and the inward diffusion of oxygen. However, the marker study does not allow any further comment on the relative diffusivities of the individual cations.
Although it might be convenient to compare the duplex oxide structure observed here with the duplex oxide structure found for TiAlN, there inevitable differences since the TiAlN/VN oxides could not be studied above around 650 °C, while the oxides on TiAlN cannot be studied <750 °C, since oxidation is so slow. Moreover, the diffusion rate of Al is low below 700 °C, while in contrast, oxygen diffuses rapidly inward through the growing oxide film and forms both Al and Ti oxides at the oxide/nitride interface. In the current study, the inner layer was TiO2, V2O5, probably Al2TiO5, and VO2, and the outer layer was Al2TiO5, V2O5, AlVO4, and TiO2. In TiAlN, a duplex oxide scale is found for temperatures of 750 °C to 900 °C, with an outer layer of Al-rich oxide and an inner layer of Ti-rich oxide.[6,7] Reduced oxidation kinetics after a temperature-dependent incubation period was associated with the passivating nature of this outer Al-rich oxide layer. However, at 900 °C, the formation of TiO2 was positively identified, which grew at an accelerated rate through cracks in the oxide layer. Interestingly, for <850 °C, the diffusion of Al and O was found to be relatively equal and noninteractive, suggesting different primary diffusion paths for the two species.
One of the drivers behind this work was to explore whether static oxidation yields the lubricious oxides of V2O5 or related Magnéli phases. While this was indeed found to be the case, the V2O5 was only dominant over a small temperature range. At 600 °C, it was only one of several phases present at the surface, with a proportion of the V being lost to AlVO4. However, a small increase in temperature to 638 °C led to a substantial change with much of the surface being covered in V2O5 crystals. However, by 674 °C, the V2O5 will have melted, although such liquids can also provide low friction. In any event, the current work confirms the rather narrow temperature range in which frictional heating would produce potentially low friction Magnéli phases at the surface.
TiAlN/VN multilayer coatings exhibited more rapid oxidation kinetics compared to TiAlN due to prevention of a passive surface oxide layer by the presence of V.
Oxidation of the TiAlN/VN coating initiated at ≥550 °C with the VN the first to oxidize. Kinetics were logarithmic initially, parabolic after 0.4 hours, and linear after ∼0.7 hours with a much slower mass gain. The behavior of the 600 °C sample was similar, except the parabolic rate was more rapid, and the oxidation rate >0.7 hours was slightly greater (0.08 g/m2 h) compared to the 550 °C sample (0.06 g/m2 h).
Cross sections by STEM/EDX and marker study revealed that the total cation outward diffusion was approximately equal to the oxygen inward diffusion. However, the diffusion rate of the V cations was clearly the most rapid as it always formed the dominant outer oxide.
At 550 °C, the oxide exhibited a rudimentary duplex oxide structure, with the outermost surface comprised of needle-/plate-shaped crystals, but the inner layer was porous and nanocrystalline. Positive phase identification could not be made, but elemental maps suggested the outermost surface was V rich while the region adjacent to the coating was dominated by Ti and Al. Phase identification of the oxide was difficult owing to the small dimensions of the particles.
At temperatures >600 °C, a duplex oxide structure was formed; the inner layer comprised a porous region of Ti-rich and V-rich nanocrystallites, while several phases were observed in the outer region, including V2O5, TiO2, and AlVO4. V2O5 was the dominant oxide at the outer layer at ≥638 °C. The outward diffusion of V depended on the species present; in the inner layer, V was present as V3+, V4+, whereas a significant V5+ was dominant in the outer layer of oxide at ≥638 °C.
We are grateful to EPSRC for funding through Grant No. GR/N23998/01.