This article appeared in a journal published by Elsevier. The attachedcopy is furnished to the author for internal non-commercial researchand education use, including for instruction at the authors institution

and sharing with colleagues.

Other uses, including reproduction and distribution, or selling orlicensing copies, or posting to personal, institutional or third party

websites are prohibited.

In most cases authors are permitted to post their version of thearticle (e.g. in Word or Tex form) to their personal website orinstitutional repository. Authors requiring further information

regarding Elsevier’s archiving and manuscript policies areencouraged to visit:

http://www.elsevier.com/copyright

Author's personal copy

Evolution of structure and properties of multi-component (AlCrTaTiZr)Ox films

Miao-I. Lin, Ming-Hung Tsai, Wan-Jui Shen, Jien-Wei Yeh ⁎Department of Materials Science and Engineering, National Tsing Hua University, 101, Section 2 Kuang Fu Road, Hsinchu 30013, Taiwan

a b s t r a c ta r t i c l e i n f o

Article history:Received 28 August 2008Received in revised form 10 October 2009Accepted 21 October 2009Available online 29 October 2009

Keywords:High-entropy alloyMulti-component oxideAmorphous structureReactive sputteringHard coating

(AlCrTaTiZr)Ox films were deposited at 350 °C by DC magnetron sputtering from high-entropy alloy target.Oxygen concentration increases with oxygen flow ratio, and saturates near 67 at.%. As-deposited films havean amorphous structure. Their hardness fall in the range of 8–13 GPa. All amorphous oxide films maintaintheir amorphous structure up to 800 °C for at least 1 h. After 900 °C 5 h annealing, crystalline phases with thestructures of ZrO2, TiO2, or Ti2ZrO6 form. Annealing enhances mechanical properties of the films. Theirhardness and modulus attain to the values about 20 and 260 GPa, respectively. The resistivity of the metallicfilms is around 102μΩcm but drastically rises to 1012μΩcm when oxygen concentration increases.

© 2009 Elsevier B.V. All rights reserved.

1. Introduction

Deposition of hard coatings is a common technology to improve theperformance of tools, dies, and molds for many different applications[1]. Binary transitionmetal nitrides such as TiN have beenwidely usedsince 1980s [2]. These simple nitrides, however, have now been sur-passed by more complex ternary systems, for example, the ternaryTi–Al–N systemwhich provides enhanced hardness and/or oxidationresistance over their binary counterparts. Exigent requirements forbetter performance push the explorations of coating materials toadvanced multi-element territory [2–4]. However, most works onprotective coatings still focus on ternary or quaternary systems.

Recently Yeh et al. start to investigate nitride coatings based onmulti-principal-element alloys, so-called “high-entropy alloys” (HEAs).Reported systems such as (AlCrTaTiZr)Nx [5–7], (AlCrSiTiV)Nx [8],(AlCrMoSiTi)Nx [9,10], and (AlMoNbSiTaTiVZr)Nx [11] display a highhardness spanning from 16 to 37 GPa. Moreover, these nitrides alsoappear to have good thermal and chemical stability [12]. Due to thepromising properties of HEA nitrides, interest also goes to their oxidecounterparts. Very recently Chen et al. reported oxide films reactivelydeposited from AlxCoCrCuFeNi (x=0.5, 1, 2) HEA targets [13]. Theirresults reveal that these oxide films also form simple solid solutionphases, and have a maximum hardness around 22.6 GPa. This is ahigh value relative to most reported oxide films, for example Al2O3

(10 GPa) [14,15], Ta2O5 (6 GPa) [16], TiO2 (18 GPa, rutile phase)[17,18], V2O5 (3–7 GPa) [19], and ZrO2 (15 GPa) [20]. One exceptionis perhaps Cr2O3, whose hardness is reported to be 25–30 GPa [21].However, several elements in their composition are not strong oxide

formers and could thus be detrimental to film properties. In this paper,we examine HEA oxides based on strong oxide-forming elements.

We selected AlCrTaTiZr as the base alloy because nitride filmsbased on this composition have been most comprehensively inves-tigated [5–7]. (AlCrTaTiZr)Nx films form one simple face-centeredcubic (FCC) NaCl-type solid-solution phase. The maximum hardnessand elastic modulus of these films without the application of bias is 32and 368 GPa, respectively. (AlCrTaTiZr)Nx films also possess goodadhesion to the substrates and good wear resistance. Here we prepare(AlCrTaTiZr)Ox films and investigate the effect of oxygen flow ratio onthe structure and properties of them. Comparisons between nitridesand oxides will also be discussed.

2. Experimental details

The AlCrTaTiZr high-entropy alloy was prepared by vacuum arcmelting. Films were deposited on Si (100) single crystal wafers andglass substrates in a DC magnetron sputtering system with a basepressure of 4×10−4Pa (3×10−6Torr). The sputtering power was150 W and substrate was kept at 350 °C during deposition. Depositionatmosphere was conducted in a mixture of Ar and O2. Chamberpressure was held constant at 0.67 Pa (5×10−3Torr) but the oxygenflow ratio, RO=O2/(O2+Ar), was varied from 0 to 50%.

The samples were annealed in a rapid thermal annealing furnacepumped to 2.4×10−4Pa (1.8×10−6Torr). The chemical compositionswere determined by electron probe X-ray micro-analysis (EPMA, JEOLJAX-8800). Both the surface morphology and film thickness of thesputtered films were observed with a scanning electron microscope(SEM, JOEL JSM-6500F FEG-SEM) operated at 15 kV. A glancing incidentangle X-ray diffractometer (GIXRD, MAC Science MXP18) was used toidentify the structure, using a CuKα radiationwith θ=1° and the2θ scan

Thin Solid Films 518 (2010) 2732–2737

⁎ Corresponding author. Tel.: +886 3 5719558; fax: +886 3 5722366.E-mail address: jwyeh@mx.nthu.edu.tw (J.-W. Yeh).

0040-6090/$ – see front matter © 2009 Elsevier B.V. All rights reserved.doi:10.1016/j.tsf.2009.10.142

Contents lists available at ScienceDirect

Thin Solid Films

j ourna l homepage: www.e lsev ie r.com/ locate / ts f

Author's personal copy

ranging from 20° to 100° at a rate of 4°min−1. The residual stress wasdetermined by Stoney's equation in which the substrate curvature wasmeasured using laser beam reflection technique. Resistivity was mea-suredbya four-point probeand thefilmhardnesswas characterizedwitha nanoindentation system (XP nanomechanical testing system, MTSCorporation) using a Berkovich indenter. The applied load for thickerfilms (RO=0 and 2.5%) is 2–5 mN, depending on the film hardness. Forthinner films (RO=15 and 50%) a smaller load range of 0.3–0.75 mN isapplied to avoid substrate effects. All the data were further confirmed tobe valid by using a silicon wafer as the calibration sample.

3. Results and discussion

3.1. Chemical composition

Fig. 1 shows the composition of the oxides films analyzed by EPMA.Oxygen concentration in the film rises sharply to 55 at.% when theoxygen flow ratio (RO) is set to 2.5%. Compared with nitride filmsprepared based on the same alloy,where the concentration of nitrogenwas only 39 at.% at amuchhigherflow ratio of 9%, it is seen that oxygenhas amuch stronger tendency to reactwith the individualmetal atomsthan nitrogen does. This is supported by the significantly larger heat offormation of themetal oxides relative to correspondingmetal nitrides.For example, the heat of formation of TiN is only −338 kJ/mol [22],while that of TiO2 is −944 kJ/mol [23]. When RO is further increased,concentration of oxygen also increases and then saturates near 67 at.%.

3.2. Structure

Fig. 2 shows the GIXRD diffraction patterns of films deposited underdifferent oxygenflowratios. All as-depositedfilmsexhibit anamorphous

structure regardless of their oxygen concentration. The amorphousstructure of the alloy film is in agreementwith our previousworks [5]. Infact, formation of amorphous alloy films have been reported in variousHEA systems including AlCrSiTiV [8], AlCrMoSiTi [9], and AlMoNbSiTa-TiVZr [11]. It is interesting to note the structure of the oxide films ofAlCrTaTiZr is also amorphous since nitride films of the same alloy havebeen reported to be crystalline and show only one single FCC NaCl-typestructure [5–7]. The discrepancy could be explained based on Table 1[22] and Table 2 [24], which list the crystal structure and lattice para-meter of the nitrides and oxides of each metallic element in the film. Itcan be seen that both the structure and lattice parameter of the nitrides,except that of AlN, are quite similar. It is, therefore, possible that thesenitrides can accommodate each other and form one structure due tohigh-entropy effect. Contrarily, a wide diversity of crystal structure andlattice parameter is seen in the oxides of the fivemetallic elements. Thatmeans if the oxide film had only one structure, for example FCC, than allothermetals except Zr has to bondwith oxygen in away totally differentto their stable configuration. This is highly unfavorable and therefore,formation of individual oxides is expected. However, during the sput-tering process there is usually not sufficient time and/or mobility for theadatoms to relocate to their equilibrium positions. The result is, there-fore, a meta-stable amorphous structure.

To justify the previous argument, annealing processes were carriedout to examine the equilibrium phases. The alloy film was excludedfrom this discussion because it reacted with the Si-substrate and lost itsintrinsic properties through the silicidation during the annealing. Incontrast, Si-substrate effect was negligible for oxide films even afterannealing at 900 °C for 5 h. From the XRD patterns, all oxide films

Fig. 1. Concentration of each element in the as-deposited films determined by EPMA asa function of oxygen flow ratio.

Fig. 2. XRD patterns of films deposited under various oxygen flow ratio.

Table 1Binary nitride characteristics of the present constituent metal elements [22].

Element Nitride Resistivity (μΩ cm) Lattice parameter (Å) Structure (RT)

Al AlN 6×1019 a: 3.1, c: 4.9 HCPCr CrN 640 4.14 FCCTa TaN 250 4.34 FCCTi TiN 100 4.24 FCCZr ZrN 300 4.58 FCC

RT: room temperature; HCP: hexagonal close-packed; FCC: face-centered cubic.

Table 2Binary oxide characteristics of the present constituent metal elements [24].

Element Oxide Band gap (eV) Lattice parameter (Å) Structure (RT)

Al Al2O3 8.8 a: 4.7, c: 12.99 HCPCr Cr2O3 4 a: 4.9, c: 13.6 HCPTa Ta2O5 4.35 – –

Ti TiO2 3 a: 4.6, c: 2.9 TetragonalZr ZrO2 5.2 5.1 FCC

RT: room temperature; HCP: hexagonal close-packed; FCC: face-centered cubic.

Fig. 3. XRD patterns of oxide films annealed at 900 °C for 5 h.

2733M.-I. Lin et al. / Thin Solid Films 518 (2010) 2732–2737

Author's personal copy

remain amorphous after annealing at 800 °C for 1 h (not shown), butcrystallizes after annealing at 900 °C for 5 h (Fig. 3). The oxide filmscrystallized into various structures such as ZrO2, TiO2, and Ti2ZrO6,which agrees with the previous explanation. Note that the thermalstability of the present amorphous (AlCrTaTiZr)Ox films is apparentlybetter than that of binary amorphous oxide films. For example, thecrystallization temperatures of amorphous Nb2O5 and Ta2O5 films are450 °Cand600–650 °C, respectively [25,26],whicharemuch lower than900 °C. The apparent discrepancy in stability between binary amor-phous oxides and the multi-principal element amorphous oxides isrelated to the high-entropy and sluggish diffusion effects in thelatter [12,27,28]. The first effect enhances the mutual solubility

among different oxides by lowering the mixing free energy of multi-oxide solution phase, thus decreases the driving force to formcrystalline phases. The second effect, originating from the difficultyfor many kinds of atoms or ions to have cooperative diffusion, canslow down the crystallization kinetics.

3.3. Morphology and deposition rate

Fig. 4 shows the plane view and cross-sectional SEM images ofoxide films deposited under different oxygen flow ratio. The metallicfilm is constructed of grains about 20 nm, and there are smallergranules in the grains. The cross-section image shows column-like

Fig. 4. FESEM images of the as-deposited films: (a) plane view and (b) cross-section images of AlCrTaTiZr metallic films; (c) plane view and (d) cross-section images of filmsdeposited at RO=2.5%; (e) plane view and (f) cross-section images of films deposited at RO=15%; (g) plane view and (h) cross-section images of films deposited at RO=50%.

2734 M.-I. Lin et al. / Thin Solid Films 518 (2010) 2732–2737

Author's personal copy

structure indicating the film growth direction. Since the XRD analysis inFig. 2 shows that the metallic film is X-ray amorphous, based on thedetailed study on amorphous film structure by Tsukimoto et al. [29], thesmaller granules are amorphous clusters, the “grains” are agglomera-tions of the clusters, and the columns are thus bundles of the ag-glomerations caused from the film growth during sputtering. Afteraddition of 2.5% oxygen to the atmosphere, the contrast of the filmsfades, the morphology becomes unclear, and the column structure

disappears indicating the film is dense with no low density columnargrain boundaries [29]. Since the XRD diffraction peak moves to smallerangle and extends to have a larger width, meanwhile there are 55 at.%oxygen in the films, the amorphous structure containing a lot of oxygenis envisaged to have a larger degree of disordering and thus smallerclusters. This amorphousfilm in fact exhibitsmetallic behavior since it isconductive aswill be proved in Section 3.5. However, the surfaces of thefilms prepared at RO=15% become rougher and show mosaic of cracknetwork. The average mosaic unit size is around 30 nm. The cross-section of the film shows lots of discrete pellets, which indicates thatfilm integrity is poor and easy to induce local cracking. The tendency togenerate cracks is reasonable since the films have a saturated content ofoxygen around 67 at.% and are actually an amorphous oxide just likeglass, as will be proved to be nonconductive in Section 3.5. At RO=50%,the morphology is similar to the mosaic of cracks at RO=15%. This isattributable to the tensile stress induced by the thermal expansioncoefficient difference between film and substrate during cooling. Theinduced tensile stress overwhelms the typical compressive intrinsicstress produced during sputtering, not only causing cracking of the filmbut also leaving someresidual tensile stress in thefilmafter cooling from350 °C to room temperature. The residual tensile stress in the films isindeed proved by Fig. 5 which shows a larger tensile stress remained inthe films deposited at higher RO.

The plane view and cross-sectional SEM images of 900 °C-annealed(AlCrTaTiZr)Ox films are shown in Fig. 6. For films with RO=2.5%, a

Fig. 5. Residual stress in the as-deposited films as a function of oxygen flow ratio.

Fig. 6. FESEM images of the oxide films after annealing at 900 °C for 5 h: (a) plane view and (b) cross-section images of films deposited at RO=2.5%; (c) plane view and (d) cross-sectionimages of films deposited at RO=15%; (e) plane view and (f) cross-section images of films deposited at RO=50%.

2735M.-I. Lin et al. / Thin Solid Films 518 (2010) 2732–2737

Author's personal copy

dense distribution of particles about 5 nm in size appear on filmsurface after annealing. These are the crystalline oxide phases formedduring annealing as seen in the XRD patterns of the annealed samplesdeposited at RO=2.5%. For films with RO of 15% and 50%, themosaic ofcracks seen in the as-deposited films has disappeared. Moreover, thecross-sectional microstructures of all annealed films also show theirintegrity consisting of dense crystalline oxide particles about 5 nm insize. However, films with RO of 15% and 50% further have a grain-likestructure with a grain size around 30 nm. This grain-like structure isinherent from the mosaic structure of cracks in the as-deposited filmsas seen in Fig. 4.

From the cross-sectional images, we can observe that film thick-ness increases slightly and then decreases dramatically to 300 nmwith increasing oxygen fraction. Fig. 7 shows the deposition rate ofthe oxide films calculated by film thickness and sputtering time. Thedrop in the deposition rate is the result of target poisoning and theresulted charging effect. In reactive sputtering, target surface mayreact with the reactive gas, forming compounds. When compoundswith low conductivity (for example, oxides) are formed, chargingeffect can occur on target surface, which significantly lowers theefficiency of the sputtering process. Moreover, oxides inherently havea lower sputtering yield relative tometals, since the covalent and ionicbonds in oxides are much stronger than metallic bonds. However,there is an increase in deposition rate at RO=2.5%. This is becausethe target poisoning is small to affect the yield rate at RO=2.5%, butthere are 55 at.% oxygen involved into the films and increases thevolume.

3.4. Hardness and Young's modulus

Film hardness and Young's modulus, both before and after an-nealing, are plotted as a function of oxygen flow ratio in Fig. 8.Hardness of the as-deposited alloy film is around 8 GPa. When theoxygen concentration of the film increases to 55 at.%, hardness israised to 13 GPa due to the incorporation of strong oxygen–metalbonding. However, hardness of the saturated oxide films decreases to8 GPa. This drop of hardness mainly resulted from the cracks in thefilms.

After annealing at 900 °C for 5h, a significant boost in film hard-ness is observed, with the highest hardness value reaching 21 GPa.This can be attributed to three enhancing factors: elimination of filmcracks, formation of nanocrystalline phases and increase in elasticmodulus. First, as film cracks in the as-deposited condition have beenremoved after annealing, it is expected hardness could be greatimproved. Second, at least three crystalline phases form afterannealing and have small grain sizes (for example, the grain size ofthe TiO2 phase at RO=2.5% is merely 14.7 nm), their significant grain-size strengthening also leads to a higher hardness [2]. Third, an

increase about 30% in Young's modulus after annealing has acorresponding contribution to hardness or strength which isinherently proportional to elastic modulus. The increase in modulusis reasonable since the as-deposited amorphous configuration ishighly non-equilibrium. During annealing atoms in the films are abletomove to their equilibrium position to construct the crystal structureand establish a larger number of bonding than that in amorphousstructure, thus increasing the modulus and hardness of the films.Although the residual stress in the film also has an effect on thehardness, it is not considered to be positive since the residual stressinduced by the difference in thermal expansion is of tensile mode, asverified by Fig. 5.

The hardness values of (AlCrTaTiZr)Ox films are, as expected,apparently lower than their nitride counterparts. However, thesevalues are still higher than most reported binary oxide films, whichrarely exceed 20 GPa (an exception is Cr2O3, which is around 25–30 GPa [21]). This indicates that these multi-element oxides can havepotential in protective coatings if the process parameters and/ortarget composition is optimized.

Fig. 9 is the hardness to elastic modulus (H/E) ratio of the filmsbefore and after annealing. It is seen that annealing increases not onlythe hardness but also the H/E ratio. The physical meaning of the H/Eratio is the absorbed energy prior to the destruction of the film. LargerH/E usually means superior wear resistance [30]. Therefore, annealinghelps the oxide films exhibit better mechanical properties.

3.5. Electrical and optical properties

From Table 2, we expect that the resistivity of AlCrTaTiZr oxidefilms should be very large. Fig. 10 displays the resistivity of the films as

Fig. 7. Deposition rate of (AlCrTaTiZr)Ox films as a function of oxygen flow ratio.

Fig. 8. Hardness and elastic modulus values of as-deposited and annealed films as afunction of oxygen flow ratio.

Fig. 9. H/E ratio of the as-deposited and annealed films as a function of oxygen flowratio.

2736 M.-I. Lin et al. / Thin Solid Films 518 (2010) 2732–2737

Author's personal copy

a function of the oxygen flow ratio. The resistivity of the HEA metallicfilm is 1.7×102μΩcm, increases to 1.4×105μΩcm at RO=2.5% andfinally to 5×1011μΩcm at even higher RO.

Films with different oxygen content exhibit different opticaltransparency. The metallic films (RO=0%) show a silvery surfaceand no optical transparency. When RO is increased to 2.5%, the filmbecome somewhat darker in tint and is still not transparent. In-creasing RO to 15% and 50% turns the films clearly transparent tonaked eye. The main optical difference between samples of RO=2.5%and 15% is not the concentration but the type of the bonding in thefilms [31]. The oxygen atoms of the sample of RO=2.5% are ininterstitial sites of solid solution. This is similar to C and N atomsoccupying the interstitial sites in a supersaturated martensite phasetransformed from austenite phase by quenching Fe alloys. For suchmetal solutions, the radii of O, C, and N are 66 pm, 77 pm, and 70 pm,respectively, based on their covalent bonds [32,33]. As the diameter ofO is larger than the hole size of octahedral site, i.e. 119.9 pm,calculated from the average atomic size of Al, Cr, Ta, Ti, and Zr, whichare 143, 128, 147, 146, and 160 pm, respectively, oxygen intersti-tials increase the interatomic spacing as reflected by the shift ofamorphous broad peak to low angle side (Fig. 2). Therefore, the filmstill presents metallic character which has a continuous bandabsorbing visible light. But, there are ionic oxide bonds in the filmof RO=15%. The band gap has formed and doesn't interact withvisible light [31]. As a result, the sample becomes transparent to thelight. The optical behavior agrees with the electrical resistivity results.

4. Conclusions

(AlCrTaTiZr)Ox films prepared by reactive sputtering are X-rayamorphous in the as-deposited state. AlCrTaTiZr film is metallicamorphous. Raising RO to 2.5% increases the concentration of oxygento 55 at.%, but the structure remains metallic amorphous. Furtherincreasing RO to 15% and 50% leads to amorphous oxides with satu-rated oxygen concentration of 67 at.%. As-deposited amorphous oxidefilms can bear 800 °C annealing for 1 h without crystallization. Crys-talline phases with structures, ZrO2, TiO2, and Ti2ZrO6 form after900 °C 5 h annealing.

The hardness of as-deposited films is around 8 GPa, but thesamples of RO=2.5% is 13 GPa resulting from the incorporation ofstrong oxygen–metal bonding. Hardness of the saturated oxide films

decreases to 8 GPa. This drop of hardness ismainly due to the cracks inthe films. Annealing greatly enhances mechanical properties by ahardness increase of 8–10 GPa and a H/E ratio increase of 0.02–0.035.The hardness increase is attributed to three enhancing factors: elimi-nation of film cracks, formation of nanocrystalline phases, and in-crease in elastic modulus.

The resistivity of the metallic films is about 102μΩcm but rapidlyraises to 1012μΩcmwhen oxygen concentration increases. The opticalbehavior agrees with the electrical resistivity results. When the re-sistivity is in metallic region, the films show a highly reflective, silverysurface. When resistivity values become oxide-like, films turns trans-parent correspondingly.

Acknowledgements

The authors gratefully acknowledge the financial supports bythe National Science Council of Taiwan under Grant No. NSC96-2628-E-007-014-MY3.

References

[1] B. North, Surf. Coat. Technol. 106 (1998) 129.[2] R. Hauert, J. Patscheider, Adv. Eng. Mater. 2 (2000) 247.[3] E. Pfluger, A. Schroer, P. Voumard, L. Donohue, W.D. Munz, Surf. Coat. Technol. 115

(1999) 17.[4] O. Knotek, W.D. Munz, T. Leyendecker, J. Vac, Sci, Technol. A—Vac. Surf. Films 5

(1987) 2173.[5] C.H. Lai, S.J. Lin, J.W. Yeh, S.Y. Chang, Surf. Coat. Technol. 201 (2006) 3275.[6] C.H. Lai, S.J. Lin, J.W. Yeh, A. Davison, J. Phys. D: Appl. Phys. 39 (2006) 4628.[7] C.H. Lai, M.H. Tsai, S.J. Lin, J.W. Yeh, Surf. Coat. Technol. 201 (2007) 6993.[8] C.H. Lin, J.G. Duh, J.W. Yeh, Surf. Coat. Technol. 201 (2007) 6304.[9] H.W. Chang, P.K. Huang, A. Davison, J.W. Yeh, C.H. Tsau, C.C. Yang, Thin Solid Films

516 (2008) 6402.[10] H.W. Chang, P.K. Huang, J.W. Yeh, A. Davison, C.H. Tsau, C.C. Yang, Surf. Coat.

Technol. 202 (2008) 3360.[11] M.H. Tsai, C.H. Lai, J.W. Yeh, J.Y. Gan, J. Phys. D: Appl. Phys. 41 (2008) 7.[12] M.H. Tsai, C.W. Wang, C.H. Lai, J.W. Yeh, J.Y. Gan, Appl. Phys. Lett. 92 (2008)

052109.[13] T.K. Chen, M.S. Wong, Thin Solid Films 516 (2007) 141.[14] T.C. Chou, T.G. Nieh, S.D. McAdams, G.M. Pharr, Scr. Metall. Materialia 25 (1991)

2203.[15] J.C. Barbour, J.A. Knapp, D.M. Follstaedt, T.M. Mayer, K.G. Minor, D.L. Linam, Nucl,

Instrum. Methods Phys. Res. Sect. B—Beam Interact. Mater. Atoms 166 (2000) 140.[16] J.E. Klemberg-Sapieha, J. Oberste-Berghaus, L. Martinu, R. Blacker, I. Stevenson, G.

Sadkhin, D. Morton, S. McEldowney, R. Klinger, P.J. Martin, N. Court, S. Dligatch, M.Gross, R.P. Netterfield, Appl. Optics 43 (2004) 2670.

[17] A. Bendavid, P.J. Martin, A. Jamting, H. Takikawa, Thin Solid Films 356 (1999) 6.[18] C.A. Freyman, Y.W. Chung, Surf. Coat. Technol. 202 (2008) 4702.[19] N. Fateh, G.A. Fontalvo, C. Mitterer, J. Phys. D—Appl. Phys. 40 (2007) 7716.[20] O. Bernard, A.M. Huntz, M. Andrieux, W. Seiler, V. Ji, S. Poissonnet, Appl. Surf. Sci.

253 (2007) 4626.[21] A.S. Kao, M.F. Doerner, V.J. Novotny, J. Appl. Phys. 66 (1989) 5315.[22] H.O. Pierson, Handbook of Refractory Carbides and Nitrides, Noyes Publications,

Westwood, New Jersey, 1996.[23] J.A. Dean (Ed.), Lange's Handbook of Chemistry, McGraw-Hill, 1999.[24] Y.M. Chiang, D.P. Birnie, W.D. Kingery, Physical Ceramics: Principles for Ceramic

Science and Engineering, Wiley-MIT, 1996.[25] J.P. Masse, H. Szymanowski, O. Zabeida, A. Amassian, J.E. Klemberg-Sapieha, L.

Martinu, Thin Solid Films 515 (2006) 1674.[26] P.H. Chang, H.Y. Liu, Thin Solid Films 258 (1995) 56.[27] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv.

Eng. Mater. 6 (2004) 299.[28] J.W. Yeh, Ann, Chim. (Paris) 31 (2006) 633.[29] S. Tsukimoto, M. Moriyama, M. Murakami, Thin Solid Films 460 (2004) 222.[30] A. Leyland, A. Matthews, Wear 246 (2000) 1.[31] D. Jiles (Ed.), Introduction to the Electronic Properties of Materials, Chapman &

Hall, London, 1994, p. 111–114, 180–196.[32] Z. Nishiyama, Martensitic Transformation, Academic Press, New York, 1978, p. 16–19.[33] C. Kittel, Introduction to Solid State Physics, John Wiley & Sons, Inc, 1996, p. 77–79.

Fig. 10. Electrical resistivity of the as-deposited films as a function of oxygen flow ratio.

2737M.-I. Lin et al. / Thin Solid Films 518 (2010) 2732–2737