Thermal stability and thermo-mechanical properties of ... - AVS

Thermal stability and thermo-mechanical properties of magnetron sputtered Cr-Al-Y-N coatings Florian Roverea兲 and Paul H. Mayrhofer Department of Physical Metallurgy and Materials Testing, Montanuniversität Leoben, A-8700 Leoben, Austria

共Received 7 August 2007; accepted 17 October 2007; published 14 December 2007兲 Cr1−xAlxN coatings are promising candidates for advanced machining and high temperature applications due to their good mechanical and thermal properties. Recently the authors have shown that reactive magnetron sputtering using Cr-Al targets with Al/Cr ratios of 1.5 and Y contents of 0, 2, 4, and 8 at % results in the formation of stoichiometric 共Cr1−xAlx兲1−yYyN films with Al/Cr ratios of ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8%, respectively. Here, the impact of Y on thermal stability, structural evolution, and thermo-mechanical properties is investigated in detail. Based on in situ stress measurements, thermal analyzing, x-ray diffraction, and transmission electron microscopy studies the authors conclude that Y effectively retards diffusional processes such as recovery, precipitation of hcp-AlN and fcc-YN, grain growth, and decomposition induced N2 release. Hence, the onset temperature of the latter shifts from ⬃1010 to 1125 ° C and the hardness after annealing at Ta = 1100 ° C increases from ⬃32 to 39 GPa with increasing YN mole fraction from 0% to 8%, respectively. © 2008 American Vacuum Society. 关DOI: 10.1116/1.2806943兴

I. INTRODUCTION Cr1−xAlxN films, where Al substitutes for Cr, have proven to be effective protective coatings for machining applications and are promising candidates for various other high temperature applications.1–7 The incorporation of Al results in a high thermal and chemical stability, allowing for high hardnesses also above deposition temperature.8 The mechanical properties 共hardness, indentation modulus, wear resistance兲 and the oxidation resistance of these coatings are scaling with the Al content. However, the mechanical properties increase only up to the metastable solubility limit of AlN in the face centered cubic 共fcc, NaCl兲 structured Cr1−xAlxN films. Exceeding this limit, which is reported to be between 70 and 80 mol %,9–12 leads to the formation of a hexagonal close packed structure 共hcp, ZnSWurtzite兲, which exhibits lower hardnesses and indentation moduli, as well as usually lower wear resistance.3,13 Consequently, the cubic modification of Cr1−xAlxN is preferred for wear resistant films in high temperature applications. When exposed to air at elevated temperatures, Cr1−xAlxN films form dense and adherent Cr2O3 / Al2O3 mixed oxides,14 explaining their superior oxidation resistance compared to other binary and ternary coating systems.15–17 Nevertheless, with industrial applications reaching or exceeding working temperatures of 1000 ° C, further optimization of the oxidation resistance and the thermal stability in general is needed as Cr1−xAlxN coatings start to release N2 at ⬃1000 ° C connected with beginning decomposition processes.8,18 This can be achieved by the incorporation of large 共substitutional兲 atoms, as they effectively retard diffusion related processes such as recovery, decomposition, and recrystallization.19–23 a兲

Electronic mail: [email protected]

29

J. Vac. Sci. Technol. A 26„1…, Jan/Feb 2008

Here especially, yttrium, being additionally a reactive element 共RE兲, most effectively also retards oxide scale growths by segregation to scale grain boundaries and the metal共coating兲/oxide interface, with the oxygen potential gradient across the scale as the driving force.24–30 While RErich segregations at the metal共coating兲/oxide interface inhibit interfacial void growth and consequently improve scale adhesion, the outward diffusion of RE ions along the scale grain boundaries effectively impedes outward cation transport, which results in retarded scale growth. In Ref. 31 we have shown that the incorporation of up to 4 at % Y into Cr1−xAlxN films with Al/Cr ratio of ⬃1.2 still allows for the formation of a single phase cubic structure. Furthermore, the Y incorporation yields increasing hardness values from ⬃31 to 38 GPa with increasing Y content from 0 to 4 at %, respectively, due to solid solution mechanisms. Here, we use simultaneous thermal analyzing 共STA兲, combining differential scanning calorimetry 共DSC兲, thermal gravimetric analysis 共TGA兲, and mass spectroscopy 共MS兲 together with in situ stress measurements during annealing, x-ray diffraction analysis 共XRD兲, nanoindentation measurements, and transmission electron microscopy 共TEM兲 to investigate the impact of Y incorporation into 共Cr1−xAlx兲1−yYyN films on their thermal stability, structure evolution, and mechanical properties for annealing temperatures 共Ta兲 up to 1500 ° C. Therefore, magnetron sputtered 共Cr1−xAlx兲1−yYyN films with Al/Cr ratios of ⬃1.2 and Y contents of 0, 1, 2, and 4 at %, corresponding to YN mole fractions of 0%, 2%, 4%, and 8%, are used. II. EXPERIMENT A laboratory scale unbalanced direct current magnetron sputtering system was used for deposition, which was carried out in an Ar+ N2 共both of 99.999% purity兲 glow discharge.

0734-2101/2008/26„1…/29/7/$23.00

©2008 American Vacuum Society

29

30

F. Rovere and P. H. Mayrhofer: Thermal stability and thermo-mechanical properties

Furthermore, powder metallurgically produced targets 共쏗75 ⫻ 6 mm, PLANSEE兲 with Y contents of 0, 2, 4, and 8 at % at a constant Al/Cr ratio of 1.5 were used. The base pressure in the system was below 1 mPa and the total working gas pressure during deposition was kept at 0.4 Pa with a N2 partial pressure of ⬃35%. The magnetron-power density was ⬃6.8 W / cm2 and the deposition temperature was set to 475 ° C. The substrates were placed parallel above the target at a distance of 50 mm. All depositions were carried out at floating potential, which was found to be −25 V for the sputtering conditions used. Nanoindentation measurements of coated MgO共100兲 substrates in the as-deposited state and after annealing in He atmosphere to temperatures up to 1200 ° C were conducted with a CSIRO ultra-micro-indentation system 共UMIS兲 using a Berkovich indentor. With respect to a proper statistic, at least 30 indents were carried out for each sample with maximum loads ranging from 15 to 45 mN. Hardnesses and indentation moduli were calculated from the loading and unloading curves after the Oliver-Pharr method.32 Phase and structural changes were monitored for postdeposition annealing temperature Ta up to 1500 ° C using simultaneously DSC, TGA, and MS in a Netzsch STA 409 and a SETSYS Evolution 共Setaram instrumentation兲 calorimeter. The measurements were carried out for a heating rate of 20 ° C / min at atmospheric pressure in flowing He 共99.999% purity, 20 sccm flow rate兲. For MS investigations 共monitoring He, N, and N2兲 the maximum temperature was limited by the equipment to 1400 ° C. To exclude substrate interference, the low alloyed steel substrates were dissolved in 10 mol % nitric acid, leaving just the coating, which was subsequently mechanically ground to a powder. Structural investigations of the as-deposited and annealed powdered films were conducted by x-ray diffraction 共XRD兲, using a Siemens D500 diffractometer in the Bragg-Brentano configuration with CuK␣ radiation 共␭ = 1.540 56 nm兲. XRD analyses of samples following STA measurements to distinct annealing temperatures and subsequent cooling down with the maximum cooling rate of 50 ° C / min 共to minimize further annealing兲 allowed for structural investigations to explain the origin of the recorded distinct endothermic or exothermic DSC features. For classification of XRD reflexes, the JCDPS database was used.33 Transmission electron microscopy 共TEM; Phillips CM12, 120 keV兲 was conducted to study the coating morphologies in the as-deposited state and after annealing in He atmosphere 共without substrate兲 to 1100 and 1200 ° C. Chemical compositions of the coatings were determined using energy dispersive x-ray analysis 共EDX; Oxford Instruments Inc.兲 in a Zeiss Evo 50 scanning electron microscope 共SEM兲. Different substrates and coating thicknesses 共t兲 were used for the individual investigations: low alloyed steel sheets for STA and x-ray diffraction, t = 9 ␮m; polished Si共100兲 plates 共20⫻ 7 ⫻ 0.35 mm3兲 for biaxial residual stress measurements, t = 1 ␮m; and polished MgO共100兲 plates 共10⫻ 10 ⫻ 1 mm3兲 for hardness measurements, t = 3 ␮m. The deposition rate of ⬃3 ␮m / h was not influenced by the target J. Vac. Sci. Technol. A, Vol. 26, No. 1, Jan/Feb 2008

30

FIG. 1. Stress-temperature cycles in vacuum up to 700 ° C of 共Cr1−xAlx兲1−yYyN films with Al/Cr ratio ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8%.

compositions used. All substrates were ultrasonically precleaned in acetone and ethanol, followed by Ar-plasma etching prior to deposition. Residual stresses ␴ and linear thermal expansion coefficients ␣ of our films were obtained using the substrate-curvature method.34 Detailed information on the measurement and the calculation of ␴ and ␣ is described in Ref. 35. III. RESULTS AND DISCUSSION Elemental analysis by SEM-EDX reveals that our 共Cr1−xAlx兲1−yYyN films are stoichiometric with N/metal ratios of 1 ± 0.02 and compositions of Cr0.46Al0.54N, Cr0.45Al0.53Y0.02N, Cr0.43Al0.53Y0.04N, and Cr0.42Al0.50Y0.08N, corresponding to Al/Cr ratios of ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8%, respectively. X-ray diffractional investigations of as-deposited samples suggest a face centered cubic structure for the whole composition range and an increasing lattice parameter from 4.11 to 4.18 Å with increasing YN mole fraction from 0% to 8%, respectively 共see also Ref. 31兲. Nanoindentation measurements of asdeposited films yield hardness values of 31.4± 2.4, 34.4± 2.5, 36.3± 1.84, and 38.1± 1.6 GPa and indentation moduli of 497± 30, 489± 26, 484± 16, and 488± 13 GPa for YN mole fractions of 0%, 2%, 4%, and 8%, respectively. Residual stress measurements indicate decreasing compressive stresses from ⬃−1.6 to −0.6 GPa with increasing YN mole fraction from 0% to 8%, respectively. In situ biaxial stress measurements during vacuum annealing to 700 ° C show linear thermo-elastic behavior up to ⬃475 ° C 共see Fig. 1兲. As soon as the deposition temperature 共⬃475 ° C兲 is exceeded, relaxation processes are activated, where growth induced point, line, and area defects change into lower energy configurations, which leads to decreasing residual compressive stresses in the films.36,37 From the slopes of the cooling segments of the stress temperature cycles and the thermal expansion coefficients ␣ of the Si substrate of 3.55⫻ 10−6 K−1 共Ref. 38兲 the ␣ values of our coatings are calculated to be ⬃8.8共±0.27兲 ⫻ 10−6 K−1. The thermal excursion up to 700 ° C results in stress relaxations

31


FIG. 2. Simultaneous thermal analyzing in inert atmosphere 共He兲 of 共Cr1−xAlx兲1−yYyN films with Al/Cr ratio ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8%: 共a兲 DSC, 共b兲 TGA, and 共c兲 MS monitoring N2.

of ⬃1.3, 1.0, 0.8, and 0.4 GPa for YN mole fractions of 0%, 2%, 4%, and 8%, respectively. Figure 2 shows simultaneous thermal analyzing 共STA兲 results of our 共Cr1−xAlx兲1−yYyN films, combining DSC 关Fig. 2共a兲兴, TGA 关Fig. 2共b兲兴, and MS monitoring N2 release 关Fig. 2共c兲兴. The base line corrected dynamical DSC signal for temperatures from 400 to 1500 ° C is composed of several overlapping exothermic and endothermic features 共indicated with A–F兲关see Fig. 2共a兲兴. Identification of the major contributions to these reactions, detected via DSC, TGA, and MS, is achieved by the combination of XRD and TEM analyses of samples annealed to various temperatures using the DSC with identical conditions. XRD measurements of powdered samples after DSC up to 870 ° C show only a small shift of the individual reflexes and therefore suggest only minute structural changes. Based on our results from in situ stress measurements 共Fig. 1兲, and as also no weight changes are detected for Ta ⱕ 1000 ° C 关see Fig. 2共b兲兴, we conclude that the exothermic features A and B 关see Fig. 2共a兲兴, over the temperature range ⌬T ⬇ 475− 900 ° C, are connected to processes involving recovery, relaxation, and beginning recrystallization, which are exothermic in nature. At Ta ⱖ 875 ° C, superimposed by feature B, a further exothermic reaction C can be detected 关see Fig. 2共a兲兴. In agreement with the literature8,18 and XRD analysis 共see Fig. 3兲, this reaction is attributed to nucleation and growth processes of hcp-AlN. The intensities of the corresponding hcpAlN reflexes at Ta = 1150 ° C are decreasing with increasing YN mole fraction, suggesting that the onset of precipitation is shifted to higher temperatures and the precipitation process is retarded. For the film containing 8 mol % YN, no hcp-AlN reflexes can be detected 共see Fig. 3兲, indicating that no hcp-AlN precipitates are formed or their volume fraction is below the ⬃2% detection limit of XRD. For the Y-free film the exothermic feature C is superimposed by a strong endothermic reaction D 共peak temperature T p ⬇ 1210 ° C兲 at Ta ⱖ 950 ° C 关see Fig. 2共a兲兴. Corresponding to findings for CrN films,39 this endothermic feature is connected with a mass loss 关Fig. 2共b兲兴 due to N2 release 关Fig. 2共c兲兴. The precipitation of hcp-AlN results in a gradually JVST A - Vacuum, Surfaces, and Films

31

FIG. 3. XRD powder scans of 共Cr1−xAlx兲1−yYyN films with Al/Cr ratio ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8% after annealing to Ta = 1150 ° C.

Cr-enriched remaining matrix and hence the retarding effect of Al on Cr-N dissociation decreases. When temperature and chemical composition provide sufficient driving force, the matrix starts to dissociate, and N2 release together with corresponding mass loss is initiated. The thereby N-depleted matrix gradually transforms to hcp-Cr2N, as indicated by XRD analysis for Ta = 1150 ° C 共Fig. 3兲. Consequently, the strong endothermic feature D 关Fig. 2共a兲兴 is attributed to Cr-N dissociation and subsequent transformation of the N-depleted matrix to hcp-Cr2N. This endothermic DSC feature 共D兲 is shifted to higher temperatures with increasing Y content 关see Fig. 2共a兲兴. The temperatures where the signals cross the zero line shift by ⬃170 ° C from ⬃1060 to 1230 ° C with increasing YN mole fraction from 0% to 8%, respectively. Consistently, also the onset temperature for mass loss connected with N2 release 关see Figs. 2共b兲 and 2共c兲兴 increases from ⬃1010 to 1125 ° C. Hence, the intensities of the hcp-Cr2N reflexes for Ta = 1150 ° C decrease with increasing YN mole fraction 共see Fig. 3兲. We therefore conclude that the large substitutional Y atoms are effective in retarding these diffusional driven processes in our 共Cr1−xAlx兲1−yYyN films. This is in excellent agreement with our results from in situ stress measurements, where the total stress relaxation by diffusional processes during annealing to 700 ° C decreases with increasing Y content. The increasing negative slope in the DSC traces at the zeroline transition to the peak of the endothermic feature D indicates an additional overlapping exothermic reaction 共E兲 in the temperature range 1150− 1250 ° C with increasing enthalpy output for increasing Y content 关see Fig. 2共a兲兴. XRD analysis of the coating with the highest Y content after annealing at Ta = 1200 ° C 共see Fig. 4兲 yields reflexes for fccYN, whereas for Ta = 1150 ° C 共see also Fig. 3兲 no YN could be detected. Consequently, the exothermic reaction E is attributed to precipitation of YN. The results obtained are in excellent agreement with selected area electron diffraction 共SAED兲 studies during plan view TEM investigations. The SAED patterns 共see Fig. 5兲 show a single phase cubic structure in the as-deposited state

32


32

FIG. 6. SAED patterns of Cr0.46Al0.54N 共0 mol % YN兲 and Cr0.42Al0.50Y0.08N 共8 mol % YN兲 after annealing to Ta = 1200 ° C with indicated diffraction rings for hcp-Cr2N, hcp-AlN, and fcc-YN. FIG. 4. XRD evolution of Cr0.42Al0.50Y0.08N 共8 mol % YN兲 with postdeposition annealing temperatures Ta up to 1500 ° C.

and also after annealing at Ta = 1100 ° C, suggesting that precipitate fractions 共hcp-AlN, hcp-Cr2N兲 are below the ⬃1% detection limit of our experimental setup, as STA 共Fig. 2兲 and XRD 共Fig. 3兲 indicate ongoing decomposition for the Y-free film. The absence of diffraction rings for hcp-Cr2N for the Y-free film can also be explained by the necessity of a considerable N depletion of the 共Cr1−xAlx兲1−yYyN matrix before, together with sufficient thermal energy, enough driving force is provided to initiate phase transformation, as is proposed by other researchers.8 The increasing diffraction ring diameters with increasing annealing temperature for Cr0.42Al0.50Y0.08N 共see Fig. 5兲 indicate a decreasing lattice

FIG. 5. SAED patterns of Cr0.46Al0.54N 共0 mol % YN兲 and Cr0.42Al0.50Y0.08N 共8 mol % YN兲 films for the as-deposited state and after annealing to Ta = 1100 ° C with indicated diffraction rings for fccCr0.46Al0.54N. J. Vac. Sci. Technol. A, Vol. 26, No. 1, Jan/Feb 2008

parameter from ⬃4.18 Å for the as-deposited state to 4.14 Å for Ta = 1100 ° C, which is in agreement with increasing diffraction angles 2␪ in the XRD patterns 共see Fig. 4兲. No comparable changes in diffraction ring diameters or diffraction angles 2␪ can be detected for Y-free Cr0.46Al0.54N, which is attributed to the overlapping of hcp-AlN precipitation 共Al reduction of the matrix兲 and N2 release, thus competitively increasing and decreasing the lattice parameter, respectively. As STA indicates no considerable N2 release for Cr0.42Al0.50Y0.08N up to Ta = 1100 ° C 关compare Figs. 2共b兲 and 2共c兲兴 and no additional phases can be detected by XRD and SAED, the decreasing lattice parameter is associated mainly with recovery processes. At Ta ⱖ 1150 ° C formation of hcp-Cr2N occurs for all coatings investigated independent from their chemical variation, whereas hcp-AlN can only be detected for the coatings with YN mole fractions of 0%, 2%, and 4% at 1150 ° C 共see Fig. 3兲. The 8 mol % YN containing film exhibits hcp-AlN formation for Ta ⱖ 1200 ° C 共see Fig. 4兲. SAED patterns after annealing at Ta = 1200 ° C 共Fig. 6兲 confirm the presence of hcp-AlN, hcp-Cr2N, and the remaining matrix for the Y-free film. In addition to these phases, fcc-YN can be detected for the film containing 8 mol % YN, in agreement with XRD analysis 共compare Fig. 4兲. These results can explain the observed morphological development of our films during annealing to 1200 ° C. The plan view TEM investigations suggest only a minute influence of Y incorporation on the coating morphology in the as-deposited state and after annealing at Ta = 1100 ° C 共compare images for Cr0.46Al0.54N and Cr0.42Al0.50Y0.08N in Fig. 7兲. Nevertheless, at temperatures above 1100 ° C, Y and the formed YN precipitates have a considerable impact on the morphological evolution of the films. After annealing at Ta = 1200 ° C the mean grain size of the film containing 8 mol % YN is much smaller compared to the Y-free film 共see Fig. 8兲. The almost dendritic morphology of the Y-free

33


FIG. 7. Plan view TEM micrographs of Cr0.46Al0.54N 共0 mol % YN兲 and Cr0.42Al0.50Y0.08N 共8 mol % YN兲 films for the as-deposited state and after annealing to Ta = 1100 ° C.

film with a mean grain size of ⬃100 nm suggests advanced decomposition, which is in good agreement with our STA and XRD results 共see Figs. 2 and 3兲 and with observations by other researchers.8 The film containing 8 mol % YN, on the other hand, exhibits a mean grain size in the range of 30 to 40 nm, which corresponds to the grain size of the asdeposited state. Although our STA investigations 关see Figs. 2共a兲–2共c兲兴 indicate that alloying of 8 mol % YN to Cr1−xAlxN most efficiently shifts the onset of N2 release to higher temperatures 共from 1010 to 1225 ° C兲, the peak temperature of the corresponding reaction D does not shift to the same extent to higher values. Films containing 4 and 8 mol % YN yield, in fact, similar peak temperatures for reaction D at ⬃1220 ° C. This is attributed to the increased grain boundary volume fraction with increasing Y content at Ta ⬎ 1100 ° C 共see Fig. 8兲, thereby providing more fast diffusion paths, together with

FIG. 8. Plan view TEM micrographs of Cr0.46Al0.54N 共0 mol % YN兲 and Cr0.42Al0.50Y0.08N 共8 mol % YN兲 films after annealing to Ta = 1200 ° C with lower and higher magnification. JVST A - Vacuum, Surfaces, and Films

33

the increased diffusivities at higher temperatures. Consistently, also the TGA signals show increasing slopes with increasing Y content in the temperature range 1150− 1250 ° C 关see Fig. 2共b兲兴. At Ta ⱖ 1250 ° C an exothermic feature 共F兲 with peak temperatures T p of 1262, 1285, 1292, and 1296 ° C is observed for our coatings with YN mole fractions of 0%, 2%, 4%, and 8%, respectively. XRD analysis confirms that this reaction is not associated with structural changes. However, as our powdered samples are compact after annealing to 1300 ° C, reaction F is assigned to processes involving sintering. Subsequent to reaction F, the DSC signal suggests another strong endothermic reaction 关indicated with G in Fig. 2共a兲兴 with peak temperatures T p of ⬃1381, 1396, 1395, and 1371 ° C for 0, 2, 4, and 8 mol % YN, respectively. This reaction is connected with a further mass loss due to N2 release 关compare Figs. 2共b兲 and 2共c兲兴. XRD investigations of our annealed samples indicate the formation of body centered cubic 共bcc兲 Cr at Ta ⱖ 1250 ° C 共see Fig. 4兲. These results suggest that the endothermic reaction G is due to decomposition of the previously formed Cr2N into Cr and N2, which is in agreement with investigations of CrN and Cr1−xAlxN films.8,18,39 With increasing Ta, the intensities of bcc-Cr reflexes continuously increase at the expense of hcpCr2N. Furthermore, the hcp-AlN and fcc-YN reflexes become more pronounced and, for Ta ⱖ 1300 ° C reflexes for the intermetallic cubic phases c-YAl, fcc-YAl2, and c-YAl3, can be detected for all Y-containing films. With YN being thermally stable up to 2670 ° C,40 the presence of these intermetallic phases is explained by formation from the N-depleted remaining matrix during sintering processes due to the formation of bcc-Cr, as observed for annealing temperatures Ta ⱖ 1300 ° C. With increasing substitutional Y content in our 共Cr1−xAlx兲1−yYyN films the heat flow intensities of the endothermic reactions D and G decrease 关see Fig. 2共a兲兴 corresponding to the decreasing amount of Cr-N bonds. Consequently, the overall mass loss due to dissociation of Cr-N to Cr and N2 becomes smaller 关see Fig. 2共b兲兴. Experimentally obtained mass losses of 11.83, 11.59, 10.81, and 10.44 wt % for YN mole fractions of 0%, 2%, 4%, and 8%, respectively, are in excellent agreement with theoretical values of ⬃12.12, 11.78, 11.10, and 10.42 wt %, obtained by subtracting the Cr-bonded N ions from Cr0.46Al0.54N, Cr0.45Al0.53Y0.02N, Cr0.43Al0.53Y0.04N, and Cr0.42Al0.50Y0.08N, respectively. Hardness values of our 共Cr1−xAlx兲1−yYyN coatings for distinct postdeposition annealing temperatures Ta are presented in Fig. 9. With beginning recovery processes, at Ta ⱖ ⬃ 475 ° C, the hardness of the Y-free film slightly decreases. The precipitation of hcp-AlN for Ta ⱖ 1000 ° C 共see previous results兲 effectively counteracts a further hardness decrease and causes a hardness maximum of 31.89± 2.09 GPa at Ta = 1100 ° C. Hence, according to STA, XRD, SAED, and TEM analysis 共Figs. 2–8兲, maximal hardening occurs at temperatures where hcp-AlN precipitates are still small and the overall film morphology is not affected by the transformation of the remaining matrix to hcp-Cr2N. For the Y-containing

34


34

tates 共in addition to hcp-AlN precipitates兲 effectively counteracts morphological changes and explains the almost constant hardness values for Ta ⱕ 1100 ° C for the Y-containing films. At Ta ⱖ 1200 ° C the ongoing decomposition of the 共Cr1−xAlx兲1−yYyN matrix to hcp-AlN, hcp-Cr2N, and fcc-YN 共for Y-containing films兲 results in a significant decrease of the film hardnesses, except for the film with the highest YN mole fraction of 8%. At Ta ⱖ 1250 ° C hcp-Cr2N further decomposes to bcc-Cr and N2 regardless of the chemical variation and yields weight losses of 11.83, 11.59, 10.81, and 10.44 wt % for YN mole fractions of 0%, 2%, 4%, and 8%, respectively. FIG. 9. Evolution of indentation hardness of 共Cr1−xAlx兲1−yYyN films with Al/Cr ratio ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8% with postdeposition annealing temperatures Ta up to 1200 ° C.

films, no considerable variation in hardness can be observed for Ta ⱕ 1100 ° C, which is in part attributed to the observed retarded diffusional processes connected with the incorporation of Y, and the precipitation of hcp-AlN and fcc-YN, thereby counteracting hardness losses due to recovery and N2 release. However, at Ta ⱖ 1100 ° C, the hardness significantly drops for all film compositions, except the 8 mol % YN containing Cr0.42Al0.50Y0.08N. While nanoindentation measurements for Ta = 1200 ° C yield hardness values of 18.18± 1.12 GPa, 22.46± 1.78 GPa, and 24.30± 1.25 GPa for mole fractions of 0%, 2%, and 4% YN, respectively, the indentation hardness for 8 mol % YN remains high at 37.54± 1.64 GPa. For the latter the retarding effect of Y incorporation on diffusional driven processes and the precipitation of hcp-AlN and fcc-YN effectively counteract morphological changes caused by decomposition induced N2 release and subsequent hcp-Cr2N formation. Consequently, at Ta = 1200 ° C, the DSC signal for YN= 8 mol % is still exothermic in nature, contrary to films with lower Y content 关see Fig. 2共a兲兴. IV. SUMMARY AND CONCLUSIONS The thermal stability and the thermo-mechanical properties of 共Cr1−xAlx兲1−yYyN films with Al/Cr ratios of ⬃1.2 and YN mole fractions of 0%, 2%, 4%, and 8% are studied. Based on STA, combining DSC, TGA, and MS monitoring N2 release together with XRD, TEM analyses, and nanoindentation measurements, we conclude that the incorporation of Y into our 共Cr1−xAlx兲1−yYyN films is effective in retarding diffusional processes. Thereby the onset of hcp-AlN precipitation and decomposition-induced N2 release with subsequent transformation of the matrix to hcp-Cr2N shift to higher temperatures. At Ta = 1200 ° C, Y-free Cr0.46Al0.54N exhibits a mean grain size of ⬃100 nm connected with advanced decomposition, whereas the 8 mol % containing Cr0.42Al0.50Y0.08N shows a mean grain size in the range of 30–40 nm, which corresponds to the grain size of the asdeposited state. Hence, the overall retardation of diffusion driven processes due to Y incorporation and fcc-YN precipiJ. Vac. Sci. Technol. A, Vol. 26, No. 1, Jan/Feb 2008

ACKNOWLEDGMENTS This work is supported by the European Commission 共project INNOVATIAL—Innovative processes and materials to synthesize knowledge-based ultra-performance nanostructured PVD thin-films on gamma titanium aluminides, Project No. NMP3-CT-2005-515884兲. 1

H. C. Barshilia, N. Selvakumar, B. Deepthi, and K. S. Rajam, Surf. Coat. Technol. 201, 2193 共2006兲. 2 M. Brizuela, A. Garcia-Luis, I. Braceras, J. I. Onate, J. C. Sanchez-Lopez, D. Martinez-Martinez, C. Lopez-Cartes, and A. Fernandez, Surf. Coat. Technol. 200, 192 共2005兲. 3 A. E. Reiter, V. H. Derflinger, B. Hanselmann, T. Bachmann, and B. Sartory, Surf. Coat. Technol. 200, 2114 共2005兲. 4 J. Romero, M. A. Gomez, J. Esteve, F. Montala, L. Carreras, M. Grifol, and A. Lousa, Thin Solid Films 515, 113 共2006兲. 5 E. Spain, J. C. Avelar-Batista, M. Letch, J. Housden, and B. Lerga, Surf. Coat. Technol. 200, 1507 共2005兲. 6 S. Ulrich, H. Holleck, J. Ye, H. Leiste, R. Loos, M. Stuber, P. Pesch, and S. Sattel, Thin Solid Films 437, 164 共2003兲. 7 J. Vetter, E. Lugscheider, and S. S. Guerreiro, Surf. Coat. Technol. 98, 1233 共1998兲. 8 H. Willmann, P. H. Mayrhofer, P. O. Å. Persson, A. E. Reiter, L. Hultman, and C. Mitterer, Scr. Mater. 54, 1847 共2006兲. 9 R. Kaindl et al., Thin Solid Films 515, 2197 共2006兲. 10 M. Kawate, J. Vac. Sci. Technol. A 20, 569 共2002兲. 11 Y. Makino and K. Nogi, Surf. Coat. Technol. 98, 1008 共1998兲. 12 A. Sugishima, H. Kajioka, and Y. Makino, Surf. Coat. Technol. 97, 590 共1997兲. 13 H. Hasegawa, M. Kawate, and T. Suzuki, Surf. Coat. Technol. 200, 2409 共2005兲. 14 O. Banakh, P. E. Schmid, R. Sanjines, and F. Levy, Surf. Coat. Technol. 163-164, 57 共2003兲. 15 W. Kalss, A. Reiter, V. Derflinger, C. Gey, and J. L. Endrino, Int. J. Refract. Met. Hard Mater. 24, 399 共2006兲. 16 M. Kawate, A. Kimura Hashimoto, and T. Suzuki, Surf. Coat. Technol. 165, 163 共2003兲. 17 R. Wuhrer and W. Y. Yeung, Scr. Mater. 50, 1461 共2004兲. 18 P. H. Mayrhofer, H. Willmann, and A. Reiter, Society of Vacuum Coaters, 49th Annual Technical Conference Proceedings 共2006兲, p. 575. 19 P. L. Gruzin, Metal Sci. Heat Treat. 2, 525 共1960兲. 20 Y. Murata, T. Kunieda, K. Yamashita, and T. Koyama, Defect Diffus. Forum 258-260, 231 共2006兲. 21 H. Ono, T. Nakano, and T. Ohta, Appl. Phys. Lett. 64, 1511 共1994兲. 22 I. Rauf and J. Boyd, Metall. Mater. Trans. A 28A, 1437 共1997兲. 23 Y. S. Sato, M. Urata, H. Kokawa, and K. Ikeda, Scr. Mater. 47, 869 共2002兲. 24 A. M. Huntz, in The Role of Active Elements in the Oxidation of High Temperature Metals and Alloys, edited by E. Lang 共Elsevier Applied Science, Amsterdam, 1989兲, p. 81. 25 B. A. Pint, Oxid. Met. 45, 1 共1996兲. 26 R. Prescott, Oxid. Met. 38, 233 共1992兲. 27 V. Provenzano, K. Sadananda, N. P. Louat, and J. R. Reed, Surf. Coat.

35


Technol. 36, 61 共1988兲. F. H. Stott, Mater. Sci. Eng. 87, 267 共1987兲. 29 J. Stringer, Mater. Sci. Eng., A 120-121, 129 共1989兲. 30 H. M. Tawancy, N. M. Abbas, and A. Bennett, Surf. Coat. Technol. 6869, 10 共1994兲. 31 F. Rovere and P. H. Mayrhofer, J. Vac. Sci. Technol. A 25, 共2007兲. 32 W. C. Oliver and G. M. Pharr, J. Mater. Res. 7, 1564 共1992兲. 33 Powder Diffraction File 共Card 35-0803 for hcp-Cr2N, Card 25-1133 for hcp-AlN, Card 35-0779 for fcc-YN, Card 06-0694 for bcc-Cr, Card 200066 for c-YAl, Card 29-0103 for fcc-YAl2, Card 20-0070 for c-Yal3兲, International Center for Diffraction Data, JCPDF-ICDD 共2005兲. 34 A. Brenner and S. Senderoff, J. Res. Natl. Inst. Stand. Technol. 42, 105 共1949兲. 28

JVST A - Vacuum, Surfaces, and Films

35

35

P. H. Mayrhofer and C. Mitterer, Surf. Coat. Technol. 133-134, 131 共2000兲. 36 P. H. Mayrhofer, A. Hörling, L. Karlsson, J. Sjölén, T. Larsson, C. Mitterer, and L. Hultman, Appl. Phys. Lett. 83, 2049 共2003兲. 37 P. H. Mayrhofer, C. Mitterer, L. Hultman, and H. Clemens, Prog. Mater. Sci. 51, 1032 共2006兲. 38 Landolt-Börnstein, Eigenschaften der Materie in ihren Aggregatzuständen, 1.Teil, Mechanisch Thermische Zustandsgrößen 共Springer, Berlin, 1971兲. 39 P. H. Mayrhofer, F. Rovere, M. Moser, C. Strondl, and R. Tietema, Scr. Mater. 57, 249 共2007兲. 40 O. N. Carlson, R. R. Lichtenberg, and J. C. Warner, J. Less-Common Met. 35, 275 共1974兲.