Surface & Coatings Technology 206 (2012) 3283–3290

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Surface & Coatings Technology

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A comparison of the oxidation behavior of CrN films deposited using continuous dc,pulsed dc and modulated pulsed power magnetron sputtering

Jianliang Lin a,⁎, Ningyi Zhang a, William D. Sproul a,b, John J. Moore a

a Advanced Coatings and Surface Engineering Laboratory (ACSEL), Department of Metallurgical and Materials Engineering, Colorado School of Mines, Golden, CO, 80401, USAb Reactive Sputtering, Inc, 2152 Goya Place, San Marcos, CA, 92078 USA

⁎ Corresponding author. Tel.: +1 303 273 3178; fax:E-mail address: jlin@mines.edu (J. Lin).

0257-8972/$ – see front matter. Published by Elsevier Bdoi:10.1016/j.surfcoat.2012.01.033

a b s t r a c t

a r t i c l e i n f o

Article history:Received 3 June 2011Accepted in revised form 16 January 2012Available online 26 January 2012

Keywords:CrN filmsModulated pulsed power (MPP) magnetronsputteringHigh power pulse magnetron sputtering(HPPMS)High power impulse power magnetronsputtering (HIPIMS)OxidationActivation energy

The study is aimed at comparing the oxidation behavior of the stoichiometric CrN films deposited by contin-uous dc magnetron sputtering (dcMS), mid-frequency pulsed dc magnetron sputtering (PMS), and modulat-ed pulsed power (MPP) magnetron sputtering techniques in a closed field unbalanced magnetron sputteringsystem. These as-deposited CrN films exhibited a cubic structure and similar stoichiometric compositions, butwith different microstructures and residual stresses. After annealing in the ambient air from 600 to 1000 °C,the changes in the crystal phase, microstructure, and hardness of the films were characterized using X-raydiffraction, scanning electron microscopy, energy dispersive X-ray spectroscopy line scan, and nanoindenta-tion. The oxidation activation energies of the films were calculated using Arrhenius equation. It was foundthat the MPP CrN film exhibited superior oxidation resistance than dcMS and PMS CrN films. After annealingat 900 °C, the MPP CrN film exhibited an extremely dense structure and the cubic phase was well maintained.On the other hand, the dcMS and PMS CrN films were severely oxidized into a porous structure with the de-velopment of β-Cr2N and Cr2O3 phases and a rapid degradation of the cubic phase after 700 and 800 °C, re-spectively. The results indicate that all films showed a parabolic oxidation rate below 900 °C. The oxidationactivation energies for the dcMS, PMS and MPP CrN films are 116 kJ/mol, 141 kJ/mol, and 195 kJ/mol, respec-tively. The better oxidation resistance of the MPP CrN film is attributed to its dense microstructure and lowresidual stress.

Published by Elsevier B.V.

1. Introduction

As one of the transition metal nitride films, CrN films have beenextensively studied for material protection due to their high hard-ness, good wear resistance, chemical inertness, and good thermal sta-bility. Magnetron sputtering is a convenient and widely usedtechnique for CrN film depositions. The CrN films have been widelydeposited using continuous dc magnetron sputtering (dcMS) [1–5],mid-frequency pulsed dc magnetron sputtering (PMS) [6] and recent-ly high power pulsed magnetron sputtering (HPPMS) techniques[7,8]. The oxidation behavior and thermal stability of the CrN filmsdeposited by dcMS and PMS have been widely reported [1–5]. In gen-eral, the cubic CrN (c-CrN) film degraded rapidly after 600 °C throughthe formation of hexagonal Cr2N (β-Cr2N) phase by the outward dif-fusion of N [9,10]. This process was accompanied by the grain growthand a decrease in the film residual stresses and defects (e.g. intersti-tials and dislocations). The observation of the Cr2O3 formation onthe CrN film surface was widely reported from 500–600 °C, whichwas dominated by the outward diffusion of Cr and inward diffusion

+1 303 273 3795.

.V.

of O through the formed Cr2O3 layer [11,12]. The oxidation rate wasfound to fit a parabolic time law.

The oxidation resistance and thermal stability of nanostructuredfilms depend onmany factors, including the film chemical composition,microstructure, residual stress, and defect density formed during thedeposition process. Djouadi et al. [13] have shown that the stoichiomet-ric c-CrN films exhibited better oxidation resistance and thermal stabil-ity than the nitrogen deficient β-Cr2N films because of the existence ofmetallic Cr in the Cr2N films. From a structural point of view, since theoxidation kinetics of a nanocrystalline film were dominated by the out-ward cationic diffusion and inward diffusion of O through the grainboundaries, the activation energy for the diffusion is largely dependenton the density, grain boundary, and grain size of the as-deposited films.It is not a surprise that the cationic and cathodic diffusions need muchhigher activation energies to pass through a dense film than in a porousfilm.

In general, the development of compressive residual stresses andthe incorporation of certain levels of defects in the hard films can ef-fectively increase the hardness of the films. However, from a thermo-dynamic point of view, these high residual stress and defect densitiesalso result in an intrinsic instability of the microstructure of the filmsat evaluated temperatures. Veprek [14] has pointed out that a num-ber of superhard films lost their superhardness whenever such films

3284 J. Lin et al. / Surface & Coatings Technology 206 (2012) 3283–3290

are annealed to a temperature higher than the deposition tempera-ture due to the relaxation of the compressive stress and defects (e.g.interstitials and dislocations) originating from the energetic ion bom-bardment during deposition. Mitterer [15] and Mayrhofer [16] sug-gested that an introduction of high defect densities to obtainsuperhard films might result in reduced thermal stability because ofthe lower thermal activation needed to start recovery in suchstructures.

Considering the above factors, one of the effective approaches toincreasing the thermal stability of nanostructured films is to increasethe density of the films while reducing the stresses and defects in thefilms by carefully controlling the deposition parameters, e.g. the ionbombardment energies and ion fluxes. Nevertheless, these plasmaconditions strongly depend on the sputtering techniques. In recentyears, HPPMS (also named high power impulse magnetron sputtering(HIPIMS)) [17–20] and its alternative technique modulated pulsedpower (MPP) magnetron sputtering [21–24] were developed to gen-erate a high degree of ionization of the target species by using pulsedhigh peak power on the target. It was suggested that the high instan-taneous deposition rate during the on time of the high power pulseleads to high initial nucleation island densities, and the intense ion ir-radiation of the growing film by the high metal ion flux providingdensification via ion bombardment-induced adatom migration [25].Ehiasarian et al. [26] improved the structure and wear resistance ofCrN films deposited by HIPIMS compared with arc bond sputtering.Recently, Lin et al. [7] also observed improvements in the density, ad-hesion, mechanical, and tribological properties of CrN films depositedby MPP compared with dcMS and PMS without using external biasand heating. Thick and dense CrN films (up to 55 μm) have been suc-cessfully deposited at a high deposition rate using the MPP technique[27].

In comparison with the traditional HPPMS/HIPIMS technique, theMPP technique has a higher deposition rate [22,28]. The MPP plasmaexhibited large metallic ion fluxes. These ions show a major peak ionenergy of 2–5 eV and a small ion energy tail with the maximum ener-gy less than 30 eV [7,23]. The large metallic ion flux in the MPP plas-ma provides enhanced ion bombardment on the growing film toachieve a fully dense structure. Nevertheless, the residual stress andthe defect density of the MPP films can be maintained low becauseof the low ion energies of the majority of the ions [7]. As discussedpreviously, these factors are critical for the thermal stability of nano-structured hard films. Therefore, it is expected that the MPP films mayexhibit a different oxidation behavior and thermal stability comparedto the films deposited by traditional magnetron sputtering techniques(e.g. dcMS and PMS). Although recent research has demonstrated theimproved microstructure, adhesion, mechanical and tribologicalproperties of several hard coating systems deposited by the HPPMS/HIPIMS and MPP techniques as compared to the traditional magne-tron sputtering techniques, no studies have been reported on thecomparison of the oxidation resistance and thermal stability ofthese films.

Therefore, the aim of this study is to use the CrN film as an exam-ple to compare the oxidation behavior of the nanostructured films

Table 1The deposition conditions for the dcMS, PMS and MPP CrN films.

Pulsing parameters Pa [kW] Pp [kW] Ia [A] Ip [A] Va [

dcMS N/A 3.0 – 7.06 – 425PMS 100 kHz and 60% duty cycle 3.0 – 8.92 – 336MPP 73 Hz and 7.3% duty cycle 3.0 95 83 190 463

• Pa and Pp are the average and peak target power.• Ia and Ip are the average and peak target current in one pulse length.• Va and Vp are the average and peak target voltage.• Id is the peak target current density during the pulse.• Isub is the substrate peak ion current density during the pulse.

deposited by dcMS, PMS, and MPP techniques with a similar stoichi-ometry and thickness, but with different microstructures and internalstresses. The CrN films were annealed from 600 °C to 1000 °C in theambient air. After the oxidation tests, changes in the crystal phase,microstructure, oxide layer thickness and hardness of the CrN filmswere characterized and compared. The oxidation activation energyof these CrN films was obtained using the Arrhenius equation.

2. Experimental details

CrN films were deposited on AISI 304 stainless steel coupons and(100) silicon wafers in a closed field unbalanced magnetron sputter-ing (CFUBMS) system by reactively sputtering a Cr target(100 mm×300 mm) in an Ar:N2 mixture. The deposition system is acylindrical chamber equipped with two unbalanced magnetrons(Teer Films, LTD.) of reversed magnetic polarities, which were placedopposite with a distance of 240 mm to form a closed magnetic field.The schematic drawing and detailed descriptions of the CFUBMS sys-tem can be found in Refs. [24,28]. The substrates were ultrasonicallycleaned in acetone and alcohol for 15 min respectively. The substrateswere installed at about 100 mm away from the Cr target. A base pres-sure below 1.3×10−4 Pa was reached prior to all depositions. Thesubstrates were sputter etched by Ar plasma bombardment at a1.34 Pa working pressure using a pulsed dc bias of −500 V(100 kHz and 90% duty cycle) for 30 min.

A thin Cr adhesion layer (100–150 nm) was deposited on the sub-strates. Then the CrN films were deposited using dcMS (Pinnacle, Ad-vanced Energy Inc.), PMS (Pinnacle Plus, Advanced Energy Inc.), orMPP (SOLO/AXIS-180™ Pulsed DC Plasma Generator, Zpulser LLC.)power sources.

Table 1 summarizes the deposition conditions. For the dcMS andPMS depositions, an average target power (Pa) of 3 kW was used. Inthe PMS depositions, an asymmetric bipolar pulsing voltage was ap-plied on the Cr target with the pulsing parameters of 100 kHz and60% duty cycle (the target voltage was reversed to 10% of the nominalnegative sputtering voltage during the positive pulse period). In theMPP depositions, the pulse length was 1000 μs, which contains a500 μs weak ionization stage and a 500 μs strong ionization stage[7]. The peak voltage (Vp), peak current (Ip) and peak power (Pp) onthe target during the strong ionization stage in the pulse were−575 V, 190 A and 95 kW, respectively, whereas Pa was maintainedconstantly at 3 kW. The repetition frequency was 73 Hz, which is re-lated to the selected Pp and Pa. A −50 V DC substrate bias voltagewas applied for all depositions (Axis, Zpulser LLC.). During the depo-sitions, high purity Ar and N2 (99.999%) were introduced into the sys-tem separately using MKS mass flow controllers. The workingpressure, which was measured by a high precision capacitance ma-nometer, was kept constant at 0.67 Pa. The total gas flow rate was62 sccm. The N2 flow rate was maintained at 50% of the total flowrate.

High temperature oxidation tests were carried out in a CARBOLITEtype 301 furnace. The samples were annealed in an ambient air at el-evated temperatures from 600 to 1000 °C for 1 h. All samples were

V] Vp [V] Id [A/cm2] Isub [mA/cm2] Pressure [Pa] Ar:N2 Bias [V]

– 0.04 1.5 0.67 1:1 −50– 0.05 4.2 0.67 1:1 −50575 1.10 215 0.67 1:1 −50

Fig. 1. Cross-sectional SEM micrographs of the as-deposited CrN films deposited using(A) dcMS, (B) PMS, and (C) MPP techniques.

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cooled down to room temperature within the furnace. Special speci-mens were made by depositing films on six faces of small pieces ofsteel substrates. The weight gain of these films was measured usingsimultaneous thermal analysis (STA) from 600 to 900 °C in the ambi-ent air for 10 h. The thermal ramping to the isothermal oxidationtemperature was set to 40 K min−1 for all samples in a 55 sccm flow-ing Ar atmosphere in order to avoid premature film oxidation. Whenthe oxidation isothermal temperature was reached, a flow of air(30 sccm) was used as the background atmosphere for the ten hourisothermal test. The oxidation rate constant (kp) was obtained fromthe weight gain as function of the time of exposure at high tempera-tures. Then the oxidation activation energy (Ea) of the films was cal-culated by plotting the inverse temperature versus the natural logof the oxidation rate using the Arrhenius equation.

ln kp� �

¼ ln Að Þ− EaR1T

where R is the gas constant, and T is the isothermal temperature.Then Ea can be obtained from the slope of the curve.

The crystal structure of the films was characterized by X-ray dif-fraction (XRD) using a Siemens diffractometer (ModelKRISTALLOFLEX-810) operated with Kα Cu radiation (20 kV and30 mA) in the conventional Bragg–Brentanomode. The residual stressof the films was measured by glancing incident angle XRD (GIXRD) inthe same X-ray diffractometer. The chemical composition of the as-deposited films was investigated by X-ray photoelectron spectrosco-py (XPS) (PHI model 5600LS) using an Al Kα X-ray source. The micro-structure of the films was characterized by a field-emission scanningelectron microscope (SEM) (JSM-7000F) using a 5 kV operation volt-age. An energy dispersive spectroscopy (EDS) line scan was carriedout along the cross-section of the annealed films to verify the oxidelayer. The hardness of the films was measured by a nanoindenter(NanoIndenter XPTM, MTS Systems Corporation) using a diamondBerkovich tip. The hardness of the annealed CrN films was measuredas a function of the indentation depth (up to 500 nm) to distinguishthe hardness of the surface oxide layer and the underlay film.

3. Results

3.1. Microstructure and properties of the as-deposited CrN films

Fig. 1 shows the cross-sectional SEM micrographs of the as-deposited CrN films. The dcMS CrN film exhibited a porous columnargrain structure, which contains long columnar grains and clear grainboundaries throughout the film thickness. The fractured morphologyof the dcMS CrN film indicates an intergranular fracture along the co-lumnar grains due to the low film density and weak boundary bond-ing strength between the grains. The PMS CrN film showed a densemicrostructure, in which fine granular grains were developedthroughout the film thickness. In contrast, the MPP CrN film exhibiteda dense and disordered short columnar grain structure. The bound-aries between the short columnar grains along the coating thicknessare not as clear as those in the dcMS film. The laminated features ob-served from the fractured cross-section of the MPP CrN film indicate atransgranular fracture across the crystals due to the high density andstrong boundary bonding strength between the grains [29]. In thisstudy, the sequence of the deposition rate from high to low is dcMS,MPP and PMS, which is similar to the results from an earlier study [7].

Fig. 2 shows the XRD patterns of the as-deposited dcMS, PMS andMPP CrN films and after annealing at different temperatures. Asshown in Fig. 2, all as-deposited films exhibit a face centered cubic(FCC) CrN (c-CrN) structure (circle symbol) (JCPDS 11–0065). ThedcMS film shows strong (111) and (200) peaks, and weak (220)and (311) peaks (Fig. 2A). The PMS film exhibits a strong (111)peak and weak (200) and (311) peaks (Fig. 2B). In contrast, the

MPP film shows a strong (311) peak and weak (111) and (200)peaks (Fig. 2C). The peaks of the PMS film shifted to a lower anglethan the dcMS and MPP films. Since the N/Cr ratios of the films aresimilar, the larger shift of the diffraction peaks in the PMS CrN filmis attributed to its higher residual stress (Table 2), which was possiblygenerated by the high energy ion bombardment from the PMS plas-ma, which contains a wide range of ion energy distribution for gasand metal ions (e.g. the maximum ion energy may reach hundredsof eV) [36]. The grain sizes of the as-deposited dcMS, PMS and MPPCrN films were 24.5 nm, 18.9 nm and 14.1 nm respectively, as calcu-lated using the Scherrer equation [30].

The thickness of the films varied in a small range of 2.8 to 3.1 μm. Assummarized in Table 1, all films have a near stoichiometric compositionwith an N/Cr ratio in the range of 0.94 to 0.98. The O content in the filmsis less than 5 at.%. The compressive residual stresses of the dcMS, PMS,and MPP CrN films are−2.0 GPa,−4.5 GPa, and−2.5 GPa, respective-ly. The hardness and H/E ratio (resistance against elastic strain to fail-ure) of the dcMS CrN film are 23 GPa and 0.07, respectively. The PMSand MPP CrN films showed similar hardness values of 28 GPa. Andthe MPP CrN film exhibited the highest H/E ratio of 0.087.

Fig. 2. XRD patterns of the as-deposited (A) dcMS, (B) PMS, and (C) MPP CrN films andafter annealing at different temperatures (the peaks indexed as ‘S’ are arisen from thesubstrate).

Table 2The properties and oxidation temperatures for the dcMS, PMS and MPP CrN films.

Technique N/Cr

Thickness[μm]

Gain size(As-dep) [nm]

Residualstress [GPa]

HardnessH [GPa]

dcMS 0.98 2.8 24.5 −2.0 23.0PMS 0.94 2.9 18.9 −4.5 27.8MPP 0.95 3.1 14.1 −2.5 28.2

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3.2. Microstructure of the films after annealing

After annealing at 600 °C, the c-CrN peaks of all three films shiftedtoward the standard positions due to the stress and defect relaxations(for interstitials and dislocations) at elevated temperatures. Strongdiffraction peaks associated with β-Cr2N phase (down triangle sym-bol) (JCPDS 35–0803) and a small amount of Cr2O3 phase (squaresymbol) (JCPDS 38–1479) were detected in the dcMS CrN film(Fig. 2A). However, nearly no crystalline oxide diffraction peakswere detected in the PMS and MPP CrN films after annealing at600 °C. The XRD patterns of the PMS film exhibited a strong c-CrNphase with the appearance of a small amount of β-Cr2N phase(Fig. 2B). In contrast, the c-CrN phase was well maintained and al-most no β-Cr2N phase was detected in the MPP CrN film after anneal-ing at 600 °C (Fig. 2C).

The cross-sectional SEM micrographs and the EDS line scan of thefilms after annealing from 700 °C to 1000 °C are shown from Figs. 3 to6, respectively.

After annealing at 700 °C, the XRD pattern of the dcMS filmshowed dominant β-Cr2N, Cr2O3, and weak Cr phases (JCPDS06–0694) (diamond label), accompanied with a rapid decrease inthe c-CrN phase (Fig. 2A). The SEM image confirms the grain growthin the top half of the dcMS film with the development of a 200 nmoxide layer on the surface (EDS line scan) (Fig. 3A). For the PMSCrN film, the intensity of the β-Cr2N (110) and (111) peaks increasedsignificantly with a rapid decrease in the c-(111) peak after annealingat 700 °C (Fig. 2B). In contrast, only a small amount of β-Cr2N phasewas formed in the MPP film (Fig. 2C). As compared to the as-deposited films, the increase in the grain size of the PMS and MPPfilms is evident from the SEM images (Fig. 3B and C). The EDS linescan shows a less than 100 nm thin oxide layer has been formed onthe PMS and MPP film surface after annealing at 700 °C, which canhardly be seen from the cross-sectional SEM images. However, thisthin oxide layer was confirmed from the XRD patterns, which showsmall Cr2O3 (012) and (104) peaks (Fig. 2B and C).

The development of the Cr2O3 phase on the CrN film surface afterannealing in the air from 600 to 700 °C has been mainly explained bythe outward diffusion of Cr species [11,12]. The phase transformationof the CrN films from c-CrN to β-Cr2N upon heat treatment as earlyas 500–600 °C has also been reported [31]. The mechanisms of thisphase transformation have been attributed to the stress relaxation[31] and outward diffusion of N in the near surface region duringannealing [10]. Therefore, the residual stress level and the density ofthe films play an important role in determining the c-CrN to β-Cr2Nphase transformation. The relatively low density of the dcMS CrN filmand the relatively high residual stress in the PMS CrN film probablyare related to their earlier phase transformation from c-CrN to β-Cr2Nat elevated temperatures than the dense MPP CrN film.

After annealing at 800 °C, the intensities of the β-Cr2N and Cr2O3

phases further increased in the dcMS CrN films, while the c-CrNphase completely disappeared (Fig. 2A). Similar to the 700 °Cannealed PMS film, the 800 °C annealed PMS film still exhibited amixture of the c-CrN and β-Cr2N phases with a slight increase in theCr2O3 peak intensities (Fig. 2B). For the MPP CrN film, small β-Cr2N(110) peak was observed with the disappearance of the c-CrN (111)peak after 800 °C (Fig. 2C). Nevertheless, the MPP CrN film still

, Young'smodulus, E [GPa]

H/E Phase formation T[°C]

StableT for cubic CrN [°C]

β-Cr2N Cr2O3

328 0.070 >600 >600 b700338 0.082 >600 >700 b800323 0.087 >700 >700 b900

Fig. 3. Cross-sectional SEMmicrographs and corresponding EDS line scans of (A) dcMS,(B) PMS, and (C) MPP CrN coatings after annealing at 700 °C. Fig. 4. Cross-sectional SEMmicrographs and corresponding EDS line scans of the (A) dcMS,

(B) PMS, and (C) MPP CrN coatings after annealing at 800 °C.

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exhibited a strong c-CrN (311) peak, which demonstrated its excel-lent thermal stability (Fig. 2C). There was a slight increase in theCr2O3 peaks of the 800 °C annealed MPP CrN film.

As shown in Fig. 4, a dense Cr2O3 layer was formed and separatedfrom the underlay large grains by a thin and porous region in thedcMS and PMS films after annealing at 800 °C. The formation of a po-rous region under the top dense oxide layer in the CrN films is mainlydue to the outward diffusion of Cr species to the surface, as widelyreported [11]. As shown in Fig. 4A and B, the top half of the dcMSand PMS CrN films exhibited a more porous structure with a largergrain size than the bottom half of the films. In contrast, the interfacebetween the top oxide layer and the MPP film after annealing at800 °C is dense. Moreover, an extremely dense microstructureunder the top oxide layer can be observed from the MPP CrN filmafter annealing at 800 °C (Fig. 4C).

After annealing at 900 °C, the intensity of the β-Cr2N peaks de-creased significantly for the dcMS CrN film, whereas the intensity ofthe Cr2O3 peaks increased continuously (Fig. 2A). The c-CrN phase inthe PMS CrN film disappeared after annealing at 900 °C. The PMS CrNfilm also showed a significant decrease in the β-Cr2N phase with a fur-ther increase in the Cr2O3 peaks (Fig. 2B). As shown in Fig. 5A and B, theSEM images and the EDS line scan confirmed that the dcMS and PMSCrN films were severely oxidized after annealing at 900 °C. The thick-ness of the oxide layer on the surface is about 600 nm and 700 nm,which can be clearly indentified from the SEM images (Fig. 5). Thestructure of the dcMS and PMS films under the top oxide layer becameextremely porous (Fig. 5A and B). The EDS line scan confirms that there

is almost no N content in the surface oxide region, and the N is reducedin the films under the oxide layer, especially for the PMS film. The largecolumnar grain boundaries in the dcMS CrN film possibly lead to thefast diffusion of N outward from the films. Although the decreasedgrain size favors the improvements in the mechanical properties, thelarge inter-crystalline volume fraction (e.g. grain boundaries) in thePMS CrN film could also act as outward diffusion paths for N atomsfrom thermodynamic point of view.

In contrast, the MPP CrN film still maintained an excellent oxida-tion resistance at 900 °C. The XRD pattern confirmed that the c-CrN(311) peak was well maintained after annealing at 900 °C (Fig. 2C).This is further justified by the SEM view of the extremely dense mi-crostructure developed under the top 600 nm dense oxide layer, asshown in Fig. 5C. The densification of the MPP CrN film at elevatedtemperatures was mainly attributed to the kinetic growth of theCrN grains with a decrease in the grain boundary areas by the energyreduction. Such a dense microstructure in the MPP film could have astrong barrier effect for slowing down the outward diffusion of theCr and N species and the inward attack from the oxygen. Consequent-ly, the MPP CrN film showed excellent phase stability.

As the oxidation temperature reached 1000 °C, Fig. 2A shows thatthe phases in the dcMS CrN films are composed of Cr2O3 and Cr, whilethe β-Cr2N phase completely disappeared. The EDS line scan detectedhigh O content and barely any N content throughout the film thick-ness, which indicates the dcMS CrN film was completely oxidized(Fig. 6A). The PMS CrN film was also severely oxidized at 1000 °C.The major phase in the PMS film was Cr2O3 with a small amount of

Fig. 6. Cross-sectional SEMmicrographs and corresponding EDS line scans of the (A) dcMS,(B) PMS, and (C) MPP CrN coatings after annealing at 1000 °C.

Fig. 5. Cross-sectional SEMmicrographs and corresponding EDS line scans of the (A) dcMS,(B) PMS, and (C) MPP CrN coatings after annealing at 900 °C.

3288 J. Lin et al. / Surface & Coatings Technology 206 (2012) 3283–3290

β-Cr2N phase (Fig. 2B). As shown in Fig. 6B, an extremely porous mi-crostructure was formed under the top fully developed oxide layer.The thickness of the complete oxidized region in the PMS CrN filmis about 2 μm accompanied with a complete reduction of N in this re-gion. However, a 1.5 μm thick region from the bottom of the PMS filmexhibited a granular grain structure without O content, which is at-tributed to the remaining β-Cr2N in the film (Fig. 6B).

In contrast, after annealing at 1000 °C, the XRD study confirmedthat the MPP CrN film contained a mixture of β-Cr2N and Cr2O3

phases (Fig. 2C). The c-CrN phase was completely transferred to β-Cr2N phase, which was now the major phase in the film. The SEMimage and EDS line scan showed that a two layered structure wasformed in the MPP CrN film, which includes a 1.5 μm polycrystallineCr2O3 layer and the underlay denseMPP CrNx film (Fig. 6C). The inter-face between the top oxide layer and the film became porous. TheEDS line scan showed that N was completely depleted in the topoxide layer, but it was maintained well in the intact film. In addition,no oxygen was detected in the underlay dense film. These results con-firmed that the inward diffusion of O and the outward diffusion of Crand N were effectively slowed down in the dense MPP CrN film.

Fig. 7. Hardness of the dcMS, PMS and MPP CrN films as a function of the oxidationtemperature.

3.3. Mechanical properties

The hardness of the dcMS, PMS and MPP CrN films was plotted asa function of the oxidation temperature as shown in Fig. 7. Afterannealing at 600 °C for 1 h, the hardness of the dcMS, PMS, andMPP CrN films dropped down to 18 GPa, 23 GPa and 25 GPa,

respectively. The decrease in the film hardness is mainly due to thestress relaxation, annealing of the defects to the grain boundaries,and also the grain growth. After 600 °C, the hardness of all films de-creased rapidly as the oxidation temperature was further increased.The hardness of the dcMS and PMS CrN films dropped down to 43%and 45% of the as-deposited hardness after annealing at 800 °C.

3289J. Lin et al. / Surface & Coatings Technology 206 (2012) 3283–3290

Nevertheless, the MPP CrN film retained 67% of the as-depositedhardness after annealing at 800 °C.

3.4. Oxidation activation energy

Plots of the square of mass gains per area (Δm)2/cm4 at differenttemperatures for the MPP CrN films with time are shown in Fig. 8A.It can be seen that the MPP CrN films exhibited a parabolic oxidationrate at the temperature range from 700 °C to 900 °C. The dcMS andPMS CrN coatings also exhibited a parabolic oxidation rate (notshown here). The parabolic oxidation rate constant (kp) at each iso-thermal oxidation temperature (T) was obtained from the slope ofthe curves using the equation:

Δmð Þ=cm2� �2 ¼ kpt:

Fig. 8B shows the Arrhenius plots of ln(kp) as a function of 1/T fordcMS, PMS and MPP CrN films. The oxidation Ea obtained from theslope of plots using the Arrhenius equation for the dcMS, PMS andMPP CrN films are 116 kJ/mol, 141 kJ/mol and 195 kJ/mol, respective-ly. These Ea values are comparable to the literature reported values forthe CrN films in a wide range of 98 to 251 kJ/mol [32–35].

Fig. 8. (A) The parabolic oxidation kinetics for the MPP CrN films at the temperaturerange from 700 °C to 900 °C, and (B) Arrhenius plot of the dcMS, PMS and MPP CrNfilms annealed at 600 °C to 900 °C in the ambient air.

4. Discussion

The above results have shown that the dcMS, PMS andMPP CrN filmsexhibited different oxidation behaviors in the ambient air. As summa-rized in Table 2, theMPP CrN film exhibits the highest stable temperaturefor c-CrN phase up to 900 °C. TheMPP CrN film also requires higher tem-peratures for the β-Cr2N phase formation as compared to the dcMS andPMS CrN films. Therefore, we conclude that the CrN films depositedusing the MPP technique exhibited superior oxidation resistance as com-pared to the films deposited by the traditional dcMS and PMS techniques.

Since the chemical composition and thickness of the CrN films aresimilar, the oxidation behavior of the CrN films is largely dependenton the microstructure, texture and residual stress of the films,which are strongly affected by the sputtering techniques.

It has been shown that the dcMS, PMS and MPP CrN depositionsgenerated different ion energy distributions and ion fluxes in theplasma [7]. The dcMS plasma exhibited low ion energies (b15 eV)and low ion fluxes; the PMS plasma exhibited a wide range of ion en-ergies (>100 eV) and a moderate level of ion fluxes; and the MPPplasma exhibited low ion energies (b30 eV) but an extensive amountof ion flux, especially for the metal ions. Therefore, the developmentof different microstructures and residual stresses of the CrN films isrelated to different plasma properties (ion energy and ion flux) ofthe dcMS, PMS and MPP plasmas. An open columnar structure wasdeveloped in the dcMS film due to insufficient ion bombardment(Fig. 1A). The high ion energies of the PMS plasma and the highmetal ion fluxes of the MPP plasma provided enhanced ion bombard-ment on the growing films, which formed a dense and less columnargrain featured microstructure. Nevertheless, the large kinetic ion en-ergies in the PMS plasma (e.g. >100 eV) also easily generated higherresidual stresses and defect densities in the PMS film [36]. In contrast,the residual stress of the MPP CrN film is lower than the PMS CrN filmbecause the ion energies of the majority of ions in the MPP plasmawere comparable to those found in the dcMS plasma.

The microstructure and residual stress differences of the CrN filmsshowed strong effects on the film oxidation resistance. The intergra-nular fracture of the dcMS CrN film is evident (Fig. 1A), which sug-gests a weak boundary bonding strength between the columnargrains. It is not surprising that the opened boundaries along thelong columnar grains can be the diffusion paths for the outward diffu-sion of the Cr and N species and the inward diffusion of O. The crystal-line Cr2O3 phase and a large amount of β-Cr2N phase were detected inthe dcMS CrN film as early as 600 °C. The c-CrN structure was onlythermally stable up to 700 °C (Fig. 2A). And the hardness of thedcMS CrN film after annealing at 800 °C maintained 43% of that ofthe as-deposited film. The oxidation activation energy of the dcMSCrN film obtained in the current study is about 116 kJ/mol.

The PMS CrN film showed the breaking down of the columnargrain growth with the development of fine granular grains but witha relatively high residual stress (−4.5 GPa) (Fig. 1B). The fine grainsize and higher residual stress of the PMS CrN film led to a high as-deposited hardness of 28 GPa. However, the large amount of grainboundaries increases the intrinsic instability of the nano-structuredstate by the occurrence of grain growth at elevated temperatures.They can also act as outward diffusion paths for N atoms from a ther-modynamic point of view. The higher residual stress of the PMS CrNfilm also results in an intrinsic instability of the microstructure ofthe films at elevated temperatures because of the lower thermal acti-vation needed to start recovery in such structures [15]. Consequently,the PMS CrN film exhibited a similar oxidation behavior to that of thedcMS CrN film. The crystalline Cr2O3 phase and a large amount of β-Cr2N phase were detected in the PMS CrN film at 700 °C. The c-CrNstructure was thermally stable up to 800 °C (Fig. 2B). And the hard-ness of the PMS CrN film after annealing at 800 °C was 45% of theas-deposited film. The oxidation activation energy of the PMS CrNfilm obtained in the current study is about 141 kJ/mol.

3290 J. Lin et al. / Surface & Coatings Technology 206 (2012) 3283–3290

The dominant transgranular fracture of the MPP CrN film suggestsa strong boundary bonding strength between the grains (Fig. 1C),which is attributed to the dense microstructure. The residual stressof the MPP CrN film is −2.5 GPa, which is comparable to that of thedcMS CrN film but lower than that of the PMS CrN film. The densestructure of the MPP CrN film can effectively reduce the inward andoutward diffusion rates of the ion species. Combined with the densestructure effect, the low residual stress in the MPP CrN film probablyslowed down the phase transformation from c-CrN to β-Cr2N. As a re-sult, the MPP CrN film exhibited superior oxidation resistance thanthe dcMS and PMS CrN films. Only a small amount of crystallineCr2O3 and β-Cr2N phases was detected in the MPP CrN film at700 °C. The c-CrN structure was thermally stable up to 900 °C(Fig. 2C). And the hardness of the MPP CrN film after annealing at800 °C was 67% of that of the as-deposited film. The MPP CrN filmexhibited an activation energy of 195 kJ/mol. It is considered thatthe higher oxidation activation energy value in the MPP CrN film isbecause of the dense microstructure which hinders the diffusion ofion species through the grain boundaries.

5. Conclusions

The oxidation behavior of the CrN films deposited by dcMS, PMSand MPP techniques was studied in the temperature range of 600 to1000 °C. The as-deposited CrN films exhibited a cubic structure andsimilar stoichiometry composition, but with different microstructuresand residual stresses. It was found that the MPP CrN film exhibitedsuperior oxidation resistance as compared to the dcMS and PMSCrN films. Only a small amount of crystalline Cr2O3 and hexagonalβ-Cr2N phases was detected in the MPP CrN film after 700 °C. Afterannealing at 900 °C, the MPP CrN film exhibited an extremely densestructure with a strong cubic (311) peak. The MPP CrN film main-tained 67% of the as-deposited hardness after annealing at 800 °C. Incontrast, the dcMS and PMS CrN films exhibited a porous structurewith a rapid degradation of the cubic phase after 700 °C and 800 °C,respectively. The results showed that all films exhibited parabolic ox-idation behavior below 900 °C. The oxidation activation energies forthe dcMS, PMS and MPP CrN films are 116 kJ/mol, 141 kJ/mol and195 kJ/mol, respectively. The improved oxidation resistance of theMPP CrN film is attributed to its dense microstructure and low resid-ual stress.

Acknowledgments

Support of this research program from the North American DieCasting Association (NADCA) and Department of Energy (DOE-OIT)is gratefully acknowledged.

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