odem

ownf thec

maiCrr aontlyosi

and the CrN/Ti6Al4V coating system have also been explored, which shows that CrN coatings can act as goodAl4V substrate.

extendue treliabioor trplicatirides,o enha

Surface & Coatings Technology 205 (2011) 46904696

Contents lists available at ScienceDirect

Surface & Coatin

l sbetter than those of TiN lm [5,6] as well as the thicker coatingforming ability of CrN due to a low compressive stress state in CrNcoating in contrast with a high compressive stress state in TiN coating[7], besides high micro-hardness and toughness [8,9]. Therefore, theuse of CrN coating on titanium alloys could be widely used to improvefriction properties and lifetime of the components in industrialapplication progressively.

Furthermore, numerous advanced surface techniques, such asnitriding [10], ion implantation [11,12], plasma spraying [13] and

nitrogen pressure on the chemical composition, structure andmechanical performance of AIP CrN coatings.

2. Experimental details

2.1. Sample preparation

All coatings were deposited in a MIP-8-800 arc ion platingsystem, using an evacuated chamber tted with a round targetphysical vapor deposition [1417], have beenenhancing the surface properties of the sphysical vapor deposition (PVD), due to its echaracteristic, convenience and precision in d

Corresponding author. Tel.: +86 24 83978081; fax:E-mail address: csun@imr.ac.cn (C. Sun).

0257-8972/$ see front matter 2011 Elsevier B.V. Adoi:10.1016/j.surfcoat.2011.04.037nce the weak surfacear resistance materialsresistance of CrN lm are

lm/substrate [20], by arc ion plating on a Ti6Al4V substrate, and havestudied not only the interfacial microstructure of the CrN/Ti6Al4Vcoating system and the tribological properties but also the effect ofperformance of the substrate as good we[3,4]. Yet, the thermal stability and corrosion1. Introduction

The Ti6Al4V alloy has been usedautomotive and biomedical industriesto-weight ratio, excellent mechanicaland biocompatibility. However, its plimited the extension of Ti6Al4V in apresistance [1,2]. Transition metal nitcoating, have usually been used tsively in the aerospace,o its attractive strength-lity, corrosion resistanceibological behavior hason areas related to wearespecially CrN and TiN

of the favorable techniques [18,19]. Nevertheless, only few researches[3] reported on the surface modication of a CrN-coated Ti6Al4V alloybymeans of an arc ion plating (AIP) process, one of the PVD processes,and neither the interface structure between the CrN layer and theTi6Al4V substrate nor their wear mechanisms have been investigatedyet.

In this study, we have deposited CrN coatings, with a Crtransitional layer in order to enhance the adhesion strength of thestudied with the aim ofubstrate. Among these,nvironmentally friendlyeposition, has been one

(diameter 64 mchromium (99.9alloy (Al: 6.02 wbalance) withthickness werewith 800-mesh(200-mesh glastially in a me

+86 24 23843436.

ll rights reserved. 2011 Elsevier B.V. All rights reserved.Ti6Al4V alloy wear resistance layer for Ti6Microstructure and mechanical propertieson Ti6Al4V substrate

Z.K. Chang, X.S. Wan, Z.L. Pei, J. Gong, C. Sun State Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Aca

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

Article history:Received 14 December 2010Accepted in revised form 5 April 2011Available online 13 April 2011

Keywords:Knoop hardness testInterfacesWearArc ion platingCrN coatings

CrN coatings have been grpressures (PN2 ). The goals ocomposition, structure and mWith an increase of PN2 , theof CrNwas transformed frommetal Cr layer, a Cr2N layeunbalanced state transitithermodynamics subsequendue to the variation in comp

j ourna l homepage: www.ef CrN coating deposited by arc ion plating

y of Sciences, 72 Wenhua Road, Shenyang 110016, PR China

by arc ion plating (AIP) onto Ti6Al4V alloy substrate at various nitrogenis investigation are to study the inuence of nitrogen pressure content on thehanical properties of AIP CrN coatings, as well as their tribological properties.n phases in the coatings changed from CrN+Cr2N+Cr to CrN, and the textureN (111)-oriented to (220)-oriented. Furthermore, themulti-layers including and a CrN layer were observed by cross-sectional TEM (XTEM), besides anlayer at the interface of CrN/substrate which was analyzed by nucleation. An increase in nitrogen pressure also resulted in a change of micro-hardnesstion and structure. Finally, the tribological properties of the Ti6Al4V substrate

gs Technology

ev ie r.com/ locate /sur fcoatm). The cathode target material was metallic% purity). Disk samples of a commercial Ti6Al4Vt.%, V: 4.10 wt.%, Fe: 0.16 wt.%, C: 0.04 wt.% and Ti:dimensions of 15 mm in diameter and 2 mm inused as the substrate. The samples were groundSiC paper and sandblasted in a wet atmospheres balls), and then ultrasonically cleaned sequen-tal detergent, acetone and deionized water,

respectively. The samples were placed on the substrate holderopposite the target surface in the vacuum chamber. The target-substrate distance was approximately 200 mm.

Prior to deposition, ion bombardment cleaning of the substrates wascarried out under900 V pulse negative bias voltage for 3 min after thebase pressure of the chamber was pumped below 7.0103 Pa. Aftercleaning, thepulsebias voltagewas reduced to150 V inorder todeposita Cr interlayer for 5 min. During the ion bombardment cleaning and Crinterlayer deposition procedures, the atmosphere of the depositionchamberwasAr gas (99.99%purity) at 0.2 Pa. ThenN2 gas (99.99%purity)

wear test parameters were as follows: a normal load of 3 N; a slidingdisplacement amplitude of 4 mm; a sliding frequency of 4 Hz and a

3. Results and discussion

3.1. Chemical composition and structure

The inuence of nitrogen pressures (PN2 ) on the chemical compo-sition of the as-deposited CrN lms was analyzed by EPMA, which wasthe result of ve different points of each sample, as shown in Fig. 1.When PN2 was 0.4 Pa, a low N content (CN33.3 at.%) was found. WithPN2 increasing over 0.8 Pa, the N content saturated at ~45 at.%. At thesame time, it can be seen that the Cr content in the CrN lms varied inthe range 6755% corresponding to the N content.

Fig. 2 shows the XRD patterns of the CrN coatings as depositedwith varying PN2 . Thesewere composed primarily of the fcc-CrN phase(JCPDS No. 76-2494) and the mixture with hexagonal-Cr2N (JCPDSNo. 35-0803) and bcc-Cr (JCPDS No. 06-0694) phases [2123]. Thecoating deposited at a lower N2 pressure (0.4 Pa) shows a (111)preferred oriented CrN phase, from which a texture developedtowards the (220) orientation as the transformation occurs withincreasing PN2 . At the same time, diffraction patterns can be analyzedusing pseudo-Voigt function prole tting [14,15,24,25], which wasinserted in Fig. 2. The intensities of the peaks corresponding to the

4691Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 46904696testing duration of 5 min. After the wear test, the wear scars of thecoatings were evaluated by an Optical Surface Proler (OSP;MicroXAM-3D, KLA-Tencor Corporation) based on the principle oflight interference, Stylus-based Surface Proler (SSP; Alpha-Step IQ,KLA-Tencor Corporation), SEM and EDS, respectively. In the stylusprolometry measurements, the scan speed, stylus force and scanlength were 50 m/s, 0.12 mN and 4 mm, respectively. The diameterof the stylus tip used in this study was 5 m.

Table 1Deposition parameters and coating thickness. The errors indicate one standard deviation.

Sample Total pressure (Pa) N2 (ml/min) Thickness (m)

No. 1 0.4 2573 3.90.3No. 2 0.8 2753 4.80.5No. 3 1.0 2824.5 5.91.0No. 4 1.2 3096 6.30.7was quickly introduced tomaintain the chamber pressure during the CrNdeposition and Ar was closed off at the same time. The depositionparameters are summarized in Table 1. A 60 A currentwas applied on theCr target during the deposition. A composite power supply (pulse biasvoltage of150 V and DC bias voltage of100 V) was employed to thesubstrates. The frequency of the pulse bias voltage was 20 kHz, and theduty cycle (ratio of the pulse duration time to a complete cycle period)was kept at 30%. The deposition time was 90 min.

2.2. Characterization of the coatings

The surface chemical compositions of the coatings were obtainedusing Electron-probe microanalysis (EPMA; EPMA-1610, Shimadzu,Japan). X-ray photoelectron spectroscopy (XPS; ESCALAB 250) wascarried out to observe the chemical bonding status in the CrN lms.The XPS spectra, obtained after removing the surface layer of samplesby sputtering with Ar+ ion for 60 s, were calibrated by carbon peak C1s at 284.5 eV. The phase structures of the coatings were character-ized by conventional BraggBrentano X-ray diffraction (XRD) using aD/max-RA type diffractometer (Cu K radiation, =1.54056 ). Themorphology and microstructure of coatings were observed by ascanning electronmicroscope (SEM; S-3000N, Hitachi, Japan) coupledwith emission dispersive spectroscopy (EDS; Oxford ISIS, UK). Thecross-sectional morphology and diffraction patterns were obtained bya Tecnai G2 F30 transmission electron microscope (TEM). Knoophardness (HK) measurements, using a load of 50 g and a dwellingtime of 15 s, were performed using an Automatic MicroindentationHardness Testing System (Model AMH43, Japan). The reciprocatingslidingwear tests were performed on a CETR UMT-2micro-tribometerunder ambient atmospheric conditions (255 C and 505% RH).During the wear tests, an actual dynamic coefcient of friction wasable to be obtained by the servo-controlled normal load. Si3N4 ballswith a diameter of 4 mm, a surface roughness Ra of 0.02 m and ahardness of HK50g 1600 were chosen as the wear counterparts. TheCr2N (111) and Cr (110) can be seen to decrease while thosecorresponding to the CrN (200) increased with increasing PN2 .

Fig. 3 shows the fraction of sub-peaks in the Cr 2p3/2 and the N 1sXPS spectra for the CrN coatings deposited at various PN2 . The spectrawere tted by the least-squares method using a GaussianLorentzianenvelope. The Cr 2p3/2 spectrum can be interpreted as beingcomposed of three species: metallic Cr0 (573.7574.4 eV [26]), Cr ina CrN environment (CrN, Cr2N, 574.5 eV [27,28]), and Cr in a CrOenvironment (e.g. Cr2O3, 575.8576.5 eV [26]). The N 1s peak wastaken to be composed of 3 groups of different chemical species: CrN(396.9 eV [28]), Cr2N (397.5 eV [27]), and the smaller peaks at399.4 0.4 eV and 401.90.4 eV which occurred in chromiumnitrites/nitrates [29]. Oxygen incorporation in the CrO [Fig. 3(a)],found in the Cr 2p3/2 spectrum, could originate from the residualoxygen gas in the chamber in which the base pressure was at a level of103 Pa [30,31]. Meanwhile, there was a much higher fraction of Cr0

(~35%) in sample 1 than that of other three samples, due toinsufcient reaction of N and Cr in the low nitrogen pressurecondition. The fraction of CrN increased quickly from 47.6% to57.9% [Fig. 3(a)] when the PN2 was increased from 0.4 to 0.8 Pa, and itwas approximately 60% between 0.8 and 1.2 Pa in the Cr 2p3/2spectrum. For the N 1s spectrum, including the CrN and Cr2Nsubpeaks, the relative concentration of CrN increased signicantly at

Fig. 1. Effects of gas condition on compositions of CrN coatings deposited on Ti6Al4V

substrate.

nitrogen pressures ranging from 0.4 to 1.0 Pa, while the Cr2Nexhibited the opposite trend over the same range, and then both ofthem changed very little [Fig. 3(b)]. These phenomena are consistentwith the EPMA results [Fig. 1], which mean that increasing the

increase in the number of CrN bonds.

Fig. 2. XRD spectra of CrN coatings deposited on a Ti6Al4V substrate at different nitrogen presCrN (220), Cr2N (111) and Cr (110).

4692 Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 46904696Fig. 3. Variations in relative intensity ratios of different chemical species of Cr and Nelements in the CrN lms affected by different nitrogen pressures: (a) CrN, Cr0 and CrOin Cr 2p3/2; (b) CrN and Cr2N in N 1s.By means of the CrN phase diagram [32], it can be concludedthat the phases present in the CrN coatings undergo a change from-Cr-Cr+-Cr2N-Cr2N-Cr2N+CrNCrN with increas-ing nitrogen content. In our experiments, the N content of sample 1 isabout 33 at.% which is in the (-Cr2N+CrN) phase region of Cr-Nphase diagram. As a result, the proportion of the Cr2N phase in sample1 is much higher than that in the other samples [Fig. 3(b)]. It is worthnoting that, although the N contents of samples 24 are all in the CrNsingle-phase region of the phase diagram, it can be still found the Cr2Nphase in the XRD patterns [Fig. 2] and the Cr2N bond in the XPSresults [Fig. 3]. Similarly, despite the fact that all four samples are inthe (-Cr2N+CrN) or CrN phase regions, there is also the Cr phase inthe XRD patterns [Fig. 2] and Cr0 bond in the XPS [Fig. 3(a)] patterns.Insufcient reaction [14] and metal droplets [14,19] during thedeposition are thought to be the reasons for the existence of the Cr2Nphase and the Cr phase, respectively. Additionally, the formation of(Cr+Cr2N) transition layers is another important factor and this willbe discussed in the next section below.

3.2. Interfacial structurenitrogen pressure results in an increase in the N content and also an

sures and the inserted image of pseudo-Voigt function prole tting of themixture withFig. 4 shows the cross-sectional TEM (XTEM) images and selected-area electron diffraction (SAED) pattern in the lm/substrateinterfacial region of the sample 4. Fig. 4(a) shows a bright eld (BF)image and a SAED pattern for the CrN coatings. These coatings exhibitmulti-layers including a metal Cr layer, a CrN layer and a Cr2N layer(the zone between thewhite dot lines andwill be proved later in Fig. 4(d)) between them. A metal Cr interlayer with a thickness of ~40 nm,formed during the pretreatment stage after the high pulse bias voltagecleaning process, and can be seen in the BF image. The correspondingdark eld (DF) image [Fig. 4(b)], obtained from the diffraction spot offcc-CrN (111) and hexagonal-Cr2N (110) which have the samecrystallographic plane distance, exhibits a strong columnar structureconsistent with the evident texture shown by XRD [Fig. 2] and SAED[Fig. 4(a)]. It can be seen from the DF image that the width of the CrNcolumns is about 3050 nm perpendicular to the interface while theCr2N layer is approximately 20 nm parallel to the interface.

Fig. 4(c) shows the BF image of CrN coatings at a highermagnication. The Cr interlayer with thickness of 3040 nm can beclearly identied. A transition layer of ~10 nm in thickness isobserved between the Cr interlayer and the Ti6Al4V substrate (thezone between the white dot lines in Fig. 4(c)). The phase in the

4693Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 46904696transition layer is not fully identied since it is too thin to becharacterized by the SAED technique alone. However, we canspeculate that it originates from the process of coating depositionby arc ion plating. Before the CrN coating depositing, the substrateneeds to endure ion bombardment with high energy in the modeof high pulse bias voltage in order to obtain a clean substrate withsome crystallographic defects (the arc cleaning process). Moreimportantly, a nanocrystalline/amorphous layer, like the transi-tion thin layer shown in Fig. 4(c), has been also found by others[18,20,3335]. Petrov et al. [18] suggested that the formation of ananocrystalline/amorphous interfacial layer might result from thehigh density of residual defect concentrations caused by the use ofhigh energy ions during the etching process.

In addition to the factors discussed above, the thermodynamicapproach based on nucleation theory is introduced to explain theformation of this unbalanced state transition layer. The reversiblework for crystal cluster formation G(r) can be expressed as a sum oftwo contributions:

G r = 4r2CV +43r3GV : 4

Where CV is the interfacial free energy per unit area between thecondensed phase and the vapor phase; where GV=(kTe/V)ln(P0/Pe)is the free energy difference per unit volume between the supersatu-rated vapor pressure Po and the equilibrium vapor pressure Pe, k theBoltzmann constant, Te the equilibrium temperature or the substrate

Fig. 4. XTEM micrograph of CrN coating deposited on Ti6Al4V substrate at 1.2 Pa N2(b) (CrNg= (111)+Cr2Ng= (110)) dark eld image; (c) highmagnication bright eld imagnetwork.temperature, V the atomvolume. The critical nucleation radius r can be

obtained by solvingG r r = 0 as follows:

r = 2CVGV

=2CVV

kTeln Po = Pe : 5

pressure: (a) bright eld image (BF) with SAED (f: fcc-CrN; h: hexagonal-Cr2N);e; and (d) HRTEM image and inset image of the lattice planes for the Cr2N interlayer

Fig. 5. Effects of gas condition on micro-hardness of CrN coatings deposited on theTi6Al4V substrate.

more obvious than the rst stage. It's attributed to that, as the testcontinues, the wear of the ball will slowly increase the contact area

4694 Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 46904696For metal Cr, the value of vapor pressure P during 2982130 K canbe obtained from the expression [36]:

lg P = kPa = 20:68 103T11:31 lg T + 13:68: 6

Using Eq. (5) in conjunction with the vapor pressure P in Eq. (6)gives the derivative of r with respect to Te:

dr = dTe =2CVV 30:19 ln Po3:02 lg Te

kT2e ln Po ln10 20:68103T1e 1:31 lg Te + 13:68 2 7

and

ln Po = ln10 20:68 103T1o 1:31 lg To + 13:68

: 8

At the same time, the equilibrium temperature Te is lower than To,which is in the level of 102103 K according to the supersaturatedvapor pressure Po. As a result, since dr /dTe0, r increases withincrease Te.

In the arc cleaning process, it can be concluded that as thesubstrate is bombarded by ions of high energy due to the high pulsebias voltage and as the substrate temperature Te rises, the criticalnucleation radius r increases. In addition, the effects of resputteringof the condensing lm and sputtering of the substrate material areenhanced with increasing bias voltage [33]. These two factors make itdifcult for the formation and growth of a stable nucleus. As a resultan unbalanced state transition layer is formed during this high

Fig. 6. Friction coefcient curve of the Ti6Al4V substrate and the CrN/Ti6Al4V coatingsystem during the ball-on-disk test.energy ion bombardment stage.In order to explore further the details of the transition layer at the

interface between the CrN and substrate, a high-resolution TEM(HRTEM) investigation has also been performed [Fig. 4(d)]. Identi-cation of the Cr2N (111) plane between the Cr layer and the CrN layerwas obtained from the lattice spacing as shown in the insetmicrograph, which has also been documented by the X-ray diffractionstudies [Fig. 2]. In a very short time after the deposit pretreatment ofthe Cr interlayer, N2 gas was gradually introduced into the vacuumchamber while Ar gas concentration was reduced to zero gradually. Athin Cr2N interlayer [Fig. 4(a) and (b)] was formed because of the lownitrogen concentration in the system at rst. On the top surface of thebcc-Cr interlayer, the hexagonal-Cr2N transition layer establishes itsepitaxial growth during the process of adding N2 and reducing the Argradually. More importantly, Cr2N phase and CrN phase have thesimilar interplanar spacing (dCr2N 110 2:40) and dCrN 111 2:39)),as shown in Fig. 2 and Fig. 4(a) and (b). So this Cr2N transition layercould reduce the internal stress in maximum, which was caused bylattice mismatch between the Cr interlayer and the fcc-CrN coating.and the friction and wear become unstable due to the accumulatedwear debris inside the groove [41].

For the CrN/Ti6Al4V system, it can be seen that the initial frictioncoefcient increased dramatically and, after some sliding, thendeclined to a stable value of 0.4, after reaching a maximum value of0.65. There are a lot of macroparticles on the surface of the CrNcoatings deposited by the AIP method and the surface of CrN/Ti6Al4Vsystem becomes rougher as a result of sandblasting. The very largemeasured friction coefcient at the beginning of the wearing stemsfrom the high local pressure caused by contact of the macroparticleson the coating surface and the Si3N4 balls which results in stressconcentration at the surrounding of the macroparticles. Thesemacroparticles were easily stripped by the load and the frictionalforce. As the result, with the prolongation of grinding-time, thesurface morphology trends to become smooth and most of macro-particles on the CrN coating were stripped and atten [Fig. 7(b) and(c)]; therefore, the friction coefcient between the coating surfaceand the Si3N4 balls decreased and stabilized gradually.

Fig. 7(a)(c) shows the SEM micrographs of the wear tracks for aTi6Al4V substrate and a CrN/Ti6Al4V coating system after 5 min offriction and wear, and Fig. 7(d) shows the local chemical elementsobtained by EDS analysis. The microhardness of the Si3N4 balls (HK14001700) is much higher than that of Ti6Al4V substrate (HK ~400).As a result, the micro-morphology of the wear cracks appears furrow-3.3. Mechanical properties

Results from Knoopmicro-hardness measurements on the Ti6Al4Vsubstrate and the CrN/Ti6Al4V coating system are shown in Fig. 5. Dueto the relatively large error, ten different points of each sample wereanalyzed. The hardness of Ti6Al4V substrate (HK ~400) was low.However, the micro-hardness is improved greatly with the CrNcoatings deposited by AIP. Additionally, with the increase of thenitrogen pressure from 0.4 to 1.2 Pa, the micro-hardness monoto-nously increases, but which uctuated substantially due to the roughsurface caused by sandblasting and some macroparticles on thecoatings. This variation may originate from three factors: rstly, solid-solution strengthening arises from the increasing of the N content[37]; secondly, the higher hardness of single phase coating than themultiphase coating [37,38] for the aforementioned reason (discussedin Section 3.1); thirdly, residual stresses [39]. In discussing the secondreason, Tian and Liu [40] suggested that besides Fermi energy theenergy of bonding and anti-bonding electron in the d-band of thematerial will be changed by composition deviation from stoichiom-etry. This could further inuence the bonding strength and result inthe phenomenon that the hardness of single phase coating is littlehigher than that of a multiphase coating. Consequently, the micro-hardness of CrN/Ti6Al4V coatings system is strongly connected withPN2 determining the structure of the coatings.

3.4. Tribological properties

To see the effect of the CrN coating on the tribological behavior ofthe Ti6Al4V substrate, tribological tests of the CrN/Ti6Al4V system(sample 2) and of the bare Ti6Al4V substrate were performed using amicro-tribometer. Fig. 6 shows the friction coefcient of the two kindsof samplesmentioned above during a ball-on-disk test. Quite differentvariations of the friction coefcient values were observed. The changeof friction coefcients for the uncoated substrate could be divided intotwo stages. In the rst stage from the beginning of wearing to 170 s,the friction (~0.35) is somewhat lower and steadier than in thesecond stage from 170 s to the end of wearing. In the second stage, thefriction coefcients show a modest gain and the uctuation range islike and the substrate suffered serious abrasive wear due to the

(a)

4695Z.K. Chang et al. / Surface & Coatings Technology 205 (2011) 46904696cutting by rough peaks on the surface of the Si3N4 balls during theprocess of friction and wear, as shown in Fig. 7(a).

The local surface morphology of the CrN/Ti6Al4V coating systemafter wear is represented in Fig. 7(b) and (c). It is noted that thesurface of wear tracks presents a scale-like morphology as well ascoating delamination and cracks [Fig. 7(c)]. During the abrasion, the

Fig. 7. SEM micrographs and EDS analysis of wear tracks under normal load of 3 N:corresponding to the point X denoted in (b).accumulation and propagation of fatigue tensile cracks caused bystress on the CrN coatings under the action of the Si3N4 balls lead tothe detachment of the coatings [41]. The CrN coatings were strippedand Si3N4 wear debris was milled into scale-like morphology. The EDSanalysis of the scale in the X-position [Fig. 7(c)] is shown in Fig. 7(d).Besides the coatings' elements, there was a relatively high content ofSi from the Si3N4 balls and O from a thin oxide tribolayer formed bythe high ash temperature during sliding [16]. No sharp groove wasobserved in the wear scar, which is mainly attributed to the fact thatthe micro-hardness of CrN coatings (HK ~1800) is higher than that ofthe Si3N4 ball. The ball-on-disk wear mechanisms of the CrN coatingson the Ti6Al4V substrate are identied as stress cracks, coatingstripping and oxidative wear.

In order to acquire the qualitative comparison of abrasion lossbetween the Ti6Al4V substrate and the CrN/Ti6Al4V coating system,an Optical Surface Proler (OSP) was used to observe the wear trackmorphologies. Fig. 8(a) and (b) displays the 3D wear trackmorphologies of Ti6Al4V substrate and CrN/Ti6Al4V coating systemrespectively. The light and deep color zones stand for higher andlower positions on the Z-axis respectively. From the wear scar prolesshown in Fig. 8(c) and (d), deep grinding cracks (~7 m) wereproduced on the Ti6Al4V substrate, whereas there was no obviouswear track in the CrN/Ti6Al4V coating system. This result demon-strates that CrN coatings are an excellent wear-resistance material toeffectively protect Ti6Al4V substrate.

4. Conclusions

We have presented experimental results and discussed themechanisms of the way which nitrogen pressure affects the chemicalcomposition, structure and mechanical performance of AIP CrNcoatings. In addition, an analysis of the interfacial microstructure ofthe lm/substrate system and comparison of the tribologicalproperties of Ti6Al4V substrates with and without coating havebeen carried out. The main results can be summarized as follows:

Ti6Al4V substrate, (b) and (c) CrN/Ti6Al4V coating system, and (d) EDS spectrum1. The N contents in CrN coating increase with increase in PN2 .Accordingly, the main phases in the lms are transformed fromCrN+Cr2N+Cr to CrN with increasing nitrogen pressure. This isbased on the phase diagram for the CrN system. And an initial(111)-dominated texture of CrN is changes into the (220) surfacesgradually.

2. Multi-layer structure of (Cr+Cr2N) interlays and CrN layer hasbeen observed by XTEM and HRTEM. In particularly, we haveobserved and theoretically proved the existence of a thinnanocrystalline/amorphous like transition layer between coatingand the substrate.

3. The micro-hardness of AIP CrN lms increases from 1550 to 2100(HK50g) as nitrogen pressure increases from 0.4 to 1.2 Pa, followingthe same trend as N and CrN contents, which is mainly ascribed tothe variation of composition and structure.

4. For CrN coatings, a combination of several wear mechanisms bystress cracks, delamination and oxidative has been consideredunder a normal load of 3 N. The application of AIP CrN coatingswith their anti-abrasive and high micro-hardness performancecould signicantly expand the application range of Ti6Al4Valloys.

Acknowledgments

The authors thank Dr. D. M. Tang, Dr. C. F. Li, Dr. B. Yang and Dr. S. J.Wang for valuable discussions about TEM during this work at IMR.Furthermore, the authors acknowledge Professor William Alan Oates(The University of Newcastle, Australia) and Dr. Z. S. You (IMR) forchecking this paper.

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substrate (a), (c) and CrN/Ti6Al4V coating system (b), (d).

Microstructure and mechanical properties of CrN coating deposited by arc ion plating on Ti6Al4V substrateIntroductionExperimental detailsSample preparationCharacterization of the coatings

Results and discussionChemical composition and structureInterfacial structureMechanical propertiesTribological properties

ConclusionsAcknowledgmentsReferences