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Available online at www.sciencedirect.com

Wear 265 (2008) 921–929

Nanoscratch and nanofriction behavior of hafnium diboride thin films

Abhishek Chatterjee a,∗, Navneet Kumar a, John R. Abelson a,Pascal Bellon a, Andreas A. Polycarpou b

a Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, 1304 W. Green Street, Urbana, IL 61801, USAb Department of Mechanical Science and Engineering, University of Illinois at Urbana-Champaign, 1206 W. Green Street, Urbana, IL 61801, USA

Received 20 February 2007; received in revised form 9 January 2008; accepted 4 February 2008Available online 2 April 2008

bstract

Hafnium diboride (HfB2) offers an excellent combination of high bulk hardness (29 GPa), high melting point (3295 ◦C) and high wear resistance,hus making it an attractive material for wear resistant coatings. In this work, the nanoscale friction response of as-deposited and annealed HfB2

lms is reported. The films were subjected to nanoscratch experimentation and the material response was investigated by measuring lateral forceso obtain friction coefficient, and by calculating Hertzian contact pressures. Both as-deposited and annealed films show favorable friction behaviorith respect to TiN. While as-deposited HfB2 and TiN deform plastically at all normal loads used in the experiments, annealed HfB2, which isanocrystalline in contrast to the X-ray amorphous as-deposited coatings, responds elastically at lower loads. The friction behavior of the films

hows no significant dependence on sliding velocity indicating negligible effect of frictional heating. The films, however, exhibit a prolongedffect of running-in which is present over the entire duration of the experiments and leads to increase in the friction coefficient with number ofasses. Annealed films display narrower and shallower scratch grooves, leading to a higher nanoscratch resistance. The excellent overall responsef annealed HfB2 films to nanoscratch testing confirms the potential of these films for wear resistant coatings.

2008 Elsevier B.V. All rights reserved.

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eywords: HfB2 coatings; Nanoscratch; Friction; Elastic–plastic response; Her

. Introduction

In machine components requiring relative motion of partsn contact, wear rates and friction coefficients generally needo be minimized [1–4]. For systems with small dimensionsnd subjected to low loads, scratching, wear and mechani-al properties on the micro- to nanoscales are very important.his is for instance the case for ultrathin films used in mag-etic storage devices [5], micro/nanoelectromechanical systemsMEMS/NEMS) [3] as well as hard coating applications [6,7].ee and Duh [6] have reported nanoscratch results of CrN films

thickness ∼ 500–1000 nm) grown by two different techniquesr.f. magnetron sputtering and cathode arc plasma). The nano-cratch resistance of these films was evaluated from the widths

nd depths of residual scratch grooves. The authors observed that.f. sputtered films exhibited higher hardness, lower root-mean-quare (rms) roughness, lower friction coefficient, and lower

∗ Corresponding author.E-mail address: achatte2@uiuc.edu (A. Chatterjee).

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043-1648/$ – see front matter © 2008 Elsevier B.V. All rights reserved.oi:10.1016/j.wear.2008.02.002

contact pressure

ear rate than plasma deposited CrN films. Another approach tovaluate the nanoscratch resistance is to explore the scratch hard-ess of the film. As part of their studies on the interdependence ofhe nanotribological performance of Ti–B–N coatings (500 nmhick) with their chemical state and structure, Ott et al. [7] had

easured a variation of the scratch hardness of the films as aunction of atomic N%. Using XPS and XRD, it was observedhat, as the N content increases, the fraction of B–N bonds andhe structural disorder increase, and the authors proposed thathis led to decreasing hardness, modulus, and scratch resistance.

In order to study the mechanisms of the friction responsexhibited by materials, it is often useful to assess its variationith normal load. In a simplified form, the coefficient of friction

an be written as the sum of plowing (μp) and adhesive (μa)omponents [8,9]. The plowing component varies linearly withormal load while the adhesive component is independent ofoad. Thus, a qualitative deduction about the mechanism can be

ade by simply observing its variation with normal load. Ottt al. [7] have used this approach to determine that the frictionesponse of Ti–B–N coatings at the nanoscale has an increasinglowing component with increasing atomic N%.

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A more specific approach is to determine whether the filmesponse to indentation and scratching is mostly elastic or plas-ic. Several studies have investigated the modes of deformationslip and densification) during nanoindentation [10]. When theanoindentation depths are extremely low, it is essential to spec-fy whether the contact areas which are used for deriving theip area function pertain to the conical region or the sphericalegion of the indenter. In a systematic study of nanoindenta-ion using both experimental and finite element analyses by Yut al. [11], the contact areas were derived for both sphericalnd conical regions of the indenter. This enabled a compari-on of the experimental contact depths with the transition depthetween spherical and conical regions of the indenter. At veryow indentation depths (spherical region), the film shows aigher component of elastic response [11].

Analysis based on the so-called Meyer’s constant, scratchardness, and Hertzian contact pressures are some of theays of studying the elastic–plastic behavior of films. Lu andomvopoulos [12] have explored the elastic–plastic transitions

n their studies of the nanomechanical and nanotribologicalroperties of amorphous carbon (a-C) films, chromium (Cr) anditanium carbide (TiC) films. They have used a phenomeno-ogical friction model proposed by Mosch et al. [13] to studyhe relative elastic–plastic behavior based on the extraction of

eyer’s constant. The model yields a power-law dependence ofhe coefficient of friction μ, with the contact load P

= kP (2−n)/n (1)

nder assumptions of homogenous materials, constant scratchelocity, and constant loading rate. The parameter n, also knowns the Meyer constant, is an indicator of the effect of contactoad on friction behavior. It has been reported that the coeffi-ient of friction of elastically deformed a-C films is inverselyroportional to the cubic root of the contact force while, forlastically deformed films, it was independent of the contactorce [14]. This suggests that n is equal to 2 and 3 for plas-ic and elastic sliding contacts, respectively. In another study15], for surfaces deformed plastically by scratching, Meyer’sonstant values 1.5 ≤ n ≤ 2 have been reported. The relativelastic–plastic response becomes very important when films areubjected to nanoindentation for very low depths (∼10 nm). Annitial indication of the mechanical response can be obtainedy comparing Hertzian contact pressure with yield strength.uch an analysis is consistent with the observation that Si(1 0 0)esponds plastically when scratched with a 100 nm diamond tiput responded elastically when the tip had a 20 �m radius [12].

However, elastic–plastic contact behavior does not providecomplete analysis of the complex phenomena involved dur-

ng nanoscratch. In particular, a simple elastic–plastic analysismodel) does not take into account parameters such as numberf passes or sliding velocity or concepts such as rolling effectf debris. As part of their studies on the tribological behavior of

LC and TiN films in the milliNewton range for different testurations, Achanta et al. [16] have reported a steady increase inhe friction coefficient of TiN and DLC films with number ofasses. They have used a concept of rolling effect of the debris

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265 (2008) 921–929

nd the source–sink effect [17] to explain their results. Theirnvestigations of the wear track using atomic force microscopyAFM) studies shed light on the importance of contact pressuresnd third body interactions in the wear track. When frictionorces vary with sliding velocity, the kinetics of the observedhenomena is of primary importance. The concept of meniscusorces is often used in nanotribodynamics and microtribodynam-cs to study changes in friction behavior with velocity [18,19].ambe and Bhushan [18] have reported on the relationshipetween the friction and adhesive forces with sliding velocitiesor Si and DLC films.

HfB2 films have been reported to have very high hardnessnd elastic modulus [20] as well as favorable wear resistance onhe macro-scale [21]. However, to the best of our knowledge,here are no reports on the nanotribological behavior of theselms, which can provide further physical insights into the behav-

or of these films. In this paper, we report on the nanofrictionehavior of the CVD-grown HfB2 films. Initial characterizationsuch as AFM, SEM, XRD and nanoindentation were performedefore the films were subjected to nanoscratch experimentationor various values of normal load, sliding velocity, and number ofasses. The material response was evaluated in terms of Hertzianontact pressure, friction coefficient and its variation with thebove parameters. Both as-deposited and annealed HfB2 filmsave been considered and compared with a high-quality TiN thinlm grown on MgO by magnetron sputtering [22].

. Experimental procedure

.1. Film deposition

HfB2 films were grown by thermal CVD of single source pre-ursor Hf(BH4)4 in a UHV chamber by the procedure describedy Jayaraman et al. [23]. The substrate used for all the growthuns was single crystal Si(0 0 1). Prior to loading in the cham-er, the substrates were degreased in acetone, isopropyl alcoholnd deionized water for 10 min each in an ultrasonic bath andnally HF dipped to remove any native oxide present. The sub-trate dimensions were 20 mm × 12 mm and they were heated byassing dc electrical current through them. The precursor pres-ure in the chamber ranged from 10−4 to 10−3 Torr. Films wererown at a substrate temperature of 260 ◦C. Some as-depositedlms were then annealed in situ at 800 ◦C for 1 h.

.2. XRD, SEM and AFM characterizations

The crystallinity of the films was analyzed by X-ray diffrac-ion. The full width at half maximum of Bragg peaks can be usedo approximate grain size by using the Scherrer formula [24]

= 0.9λ

B cos θ(2)

here t is grain size, B is the broadening factor, λ is the wave-

ength of X-rays (1.54 A), and θ is the angle of Bragg peakeflections. The measured full-width at half-maximum (Bm) hado be corrected for instrumental broadening (Bi = 0.15◦) usinghe equation: B2 = B2

m–B2i [24].

Wear 265 (2008) 921–929 923

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Table 1Summary of experimental parameters for nanoscratch experiments

Coatings As-deposited HfB2, annealed HfB2, TiN

Normal loads (�N) 100, 200, 350, 500SN

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A. Chatterjee et al. /

Cross-sectional SEM of the films was performed to studyhe film morphology. The roughness of the films was mea-ured using a tapping mode AFM which is part of a MultimodeFM (Nanoscope E, Digital Instruments). Surface areas of�m × 3 �m were imaged using a 10 nm sharp Si tip. The

ms roughness was determined from ten measurements obtainedrom different areas of each film.

.3. Nanoindentation

Film hardness (H) and reduced elastic modulus (Er) wereeasured by nanoindentation using the procedure proposed byliver and Pharr [25]. A Berkovich indenter was used for the

ests and the area functions used for calculation of H and Er wereerived by fitting nanoindentation data on standard fused quartzample to the function

= ABer + C1hc + C2h1/2c + C3h

1/4c + C4h

1/8c + C5h

1/16c

(3)

here hc represents the contact depth and

Ber = 24.5h2c (4)

s the area function for a perfect Berkovich tip. As shown in Fig.in Yu et al. [11], the Berkovich tip could be approximated asconical indenter with a spherical tip. The semi-vertical angle

or an ideal Berkovich indenter is θ = 70.3◦. The area functionsor the conical and spherical parts are thus given by

sph = πhc(2R − hc), for hc ≤ ha (5)

con = π tan2 θ(hc + h0)2, for hc > ha (6)

nd the change from the spherical to the conical region occurst

a = R − R sin θ (7)

The radius of the Berkovich tip used for the nanoindenta-ion experiments was inferred from shallow nanoindentationxperiments on a standard quartz sample provided by the man-facturer (Hysitron Inc.). The obtained tip radius, ≈150 nm, isn agreement with direct measurements obtained from high res-lution SEM images of the Berkovich tip, performed in thistudy. The elastic moduli (Ef) of the films were obtained fromhe expression for the reduced modulus

1

Er= 1 − ν2

d

Ed+ 1 − ν2

f

Ef(8)

Poisson ratios for diamond (νd = 0.07), film (νf = 0.12) [26]nd the elastic modulus of diamond (Ed = 1141 GPa) were usedor the calculations. A typical loading profile consisted of mul-iple loading–unloading cycles on exactly the same location

termed “pull” loading) with successively increasing loads tomaximum of 1 mN. The reported hardness values correspond

o an indenter contact depth less than 10% of the film thickness,hus avoiding substrate effects.

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liding velocities (�m/s) 0.32, 0.44, 0.64, 0.87, 1.3o. of passes 1, 3, 5, 10, 20, 50

.4. Nanoscratch

Nanotribological experiments were performed by means ofnanoscratch tester (Triboscope 2D, Hysitron Inc.) interfacedith a multimode atomic force microscope (Nanoscope E, Dig-

tal Instruments). To eliminate the effect of tip geometry oncratching the film, a conospherical shaped diamond probe ofadius 870 nm (as measured from SEM images) was used toonduct the nanoscratch tests. These tests were performed withonstant normal loads of 100, 200, 350 and 500 �N. Due to sys-em constraints, loads above 500 �N were not used. All scratchengths were kept constant at 6 �m. The normal force, lateralfriction) force, normal displacement and lateral displacementuring scratch tests were recorded with respect to time. Coeffi-ients of friction (the ratio of the measured friction force to thepplied normal force) were calculated as an average of about00 data points over the entire sliding distance. To study theffect of frictional heating, the sliding velocities were variedn the range of 0.3–1.3 �m/s. The number of passes was var-ed from 1 to 50 in order to investigate running-in transientsnd their effect on the friction behavior. The mechanical driftue to small imprecisions during the repositioning of the tip asell as thermal drift places a limit on the maximum number ofasses, about 50 with the present system. AFM impressions ofanoscratch grooves were obtained with the conospherical dia-ond tip to measure the residual depths and scratch widths. Thisas performed using the SPM mode of the instrument where the

ame region was scanned at much reduced loads (∼0.9 �N). Thebove nanoscratch experiments were also performed on a stan-ard TiN thin film for comparison. The experimental conditionsnd test parameters are summarized in Table 1.

. Results

.1. SEM, AFM and XRD of both as-deposited andnnealed films

Cross-sectional SEM images of the aforementioned filmsFig. 1a and b) show that the films are dense columnar andmooth. The thickness of as-deposited films is estimated to be220 nm while that of annealed films is ∼150 nm. The thickness

hange is due to densification during annealing. AFM imagesFig. 2a and b) show that asperities are distributed over a broaderange of heights for as-deposited films compared to annealedlms. This is confirmed by the roughness values, Rrms of 1.8

nd 1.4 nm for as-deposited and annealed films, respectively.oth the AFM images and the Rrms values prove that anneal-

ng reduces the overall roughness of the films. XRD spectra ofhe as-deposited films show very broad Bragg diffraction peaks

924 A. Chatterjee et al. / Wear 265 (2008) 921–929

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ig. 1. Cross-sectional SEM images of (a) as-deposited HfB2 film and (b)nnealed HfB2 film.

hile the annealed sample shows well-defined broad Braggeaks (Fig. 3a and b). The peaks recorded for as-grown films arendicative of the presence of very small fraction of crystallites inlargely amorphous matrix. The apparent grain size calculatedsing the Scherrer equation is 1.3 nm. For the annealed films,he calculated grain size is 6.5 nm. These results are consistentith TEM characterization of similar HfB2 films [20].

.2. Nanoindentation

Annealed HfB2 films reach the highest hardness of all the

lms at 43 GPa, which is similar to the value obtained by Jayara-an et al. [20] (40 GPa). TiN and as-deposited HfB2 films show

ardness values of 15 and 22 GPa, respectively. These resultsre summarized in Table 2. According to Eq. (7) and setting

able 2ilm parameters measured prior to nanotribological testing

oatings Thickness (nm) Rrms (nm) H (GPa) Ef (GPa)

s-deposited HfB2 220 1.8 22 378nnealed HfB2 150 1.4 43 473iN 300 2.1 15 410

F

ig. 2. AFM images of (a) as-deposited HfB2 film and (b) annealed HfB2 film.

= 150 nm, the depth of the spherical part ha, is 8.78 nm. Theaximum indentation depth (during the last ramp step of the

ull-loading profile in our experiments was 15 nm (7.5% oflm thickness) for as-deposited films and 10 nm (6.7% of film

hickness) for annealed films. The corresponding contact depthshc) are 11 and 8 nm for as-deposited films and annealed films,espectively. Therefore, for as-deposited films, the contact areasre beyond the spherical region and pertain only to the conical

ig. 3. XRD plots of (a) as-deposited HfB2 film and (b) annealed HfB2 film.

A. Chatterjee et al. / Wear

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ig. 4. Variation of the Hertzian contact pressures with normal load, and com-arison with the film yield strengths.

.3. Nanoscratch

.3.1. Hertzian contact pressureFig. 4 depicts the variation of the calculated Hertzian con-

act pressures at the onset of the nanoscratch experiments andcomparison with films’ yield strengths. The yield strengthsere estimated from Tabor’s equation, H = 2.7Y, where H is

he hardness obtained from nanoindentation experiments andis the yield strength. For as-deposited HfB2 and TiN films,

he Hertzian contact pressure exceeds the yield strengths at allpplied normal loads. For annealed HfB2, however, the Hertzianontact pressure is lower than the yield strength for loads below200 �N. Thus we expect that as-deposited HfB2 and TiN filmsill deform plastically for the whole range of loads used in this

tudy. Annealed HfB2, in contrast, should respond elastically, ateast for the 100 and 200 �N normal load experiments.

.3.2. Load-dependence of frictionIn the first set of nanoscratch experiments, the friction behav-

or was studied at different normal loads (100–500 �N) keepinghe sliding velocity and the number of passes constant. AFMurface morphologies of the residual scratches are shown inig. 5a–c for as-deposited HfB2, annealed HfB2 and TiN films,espectively. Annealed HfB2 shows shallower and narrowercratches compared to as-deposited HfB2. The variations of fric-ion coefficient with normal load are shown in Fig. 6, whererror bars designate one standard deviation, obtained from typi-ally 10 independent measurements. The friction coefficients ofiboride films, 0.09–0.11 for as-deposited HfB2 and 0.08–0.09or annealed HfB2, are significantly lower than that of TiN0.14–0.17). μ increases with increasing contact load for as-eposited HfB2and TiN film. The value of n (Meyer’s constant)xtracted from these results is 1.78 for TiN and 1.71 for as-

eposited HfB2 films. On the other hand, the coefficient ofriction decreases with normal load for annealed HfB2 films,nd this is reflected in the Meyer’s constant value of 2.43. ForiN and as-deposited HfB2 films, the increase in friction force

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265 (2008) 921–929 925

reflected in the increase in μ) with normal load is a manifes-ation of the larger plowing friction component. For annealedlms, the adhesive component dominates over a smaller plowingomponent. The nanoscratch resistance of each surface can bevaluated qualitatively by observing the differences in the resid-al scratch morphology (Fig. 5). It is clear that annealed HfB2ffer higher nanoscratch resistance compared to as-depositedfB2 and TiN. Annealing improves the friction behavior of HfB2

ppreciably and this effect becomes more pronounced as the loads increased.

.3.3. Dependence on sliding velocityFurther insights can be obtained by studying the variation

f friction coefficient with respect to sliding velocity where theormal load and number of passes were kept constant. Fig. 7epicts the AFM image of residual scratches obtained in TiNor the aforementioned set of experiments. Fig. 8 depicts theariation of μ when the sliding velocities were varied in theange of 0.3–1.3 �m/s at a constant normal load of 200 �N.

None of the films show any appreciable dependence of theriction coefficient with sliding velocity; it can thus be inferredhat frictional heating is not sufficiently strong to make its pres-nce felt on the overall friction data. Even if the interface flashemperature may change at the tips of some of the asperities,he overall heat generated is negligible and dissipated in thenderlying material.

.3.4. Dependence on the number of passesFig. 9 depicts the friction coefficient variation with the num-

er of passes for the three films. The trends are similar for allhe films as all of them show a steady increase in μ with increas-ng cycle number. We note however that for TiN, a quasi steadytate appears to have been reached. Fig. 10 depicts the variationf contact depths and contact areas with the number of passesor both as-deposited and annealed HfB2 films at normal loadf 500 �N, and sliding velocity of 0.32. Both contact depthsnd contact areas increase at a much faster rate for as-depositedfB2 compared to annealed HfB2. Fig. 11 compares the per-

entage change in contact area (�A) versus number of passes tohe percentage change in friction coefficient (�μ) versus num-er of passes. It can be observed that �A and �μ are similaror as-deposited HfB2 at the same number of passes. �μ at 50asses is 63.6% while �A at 50 passes is ∼50%. On the otherand, annealed HfB2 and TiN (not shown in the figure) do nothow any correlation between �A and �μ. For example, fornnealed HfB2, �μ (87.5% at 50 passes) is much greater thanA (11.5% at 50 passes).

. Discussion

In the nanoindentation experiments, the contact depths (hc)or as-deposited and annealed HfB2 films have a significantffect on their relative elastic–plastic response. For as-deposited

lms, hc (11 nm) is greater than the spherical–conical transitionepth ha. It can thus be safely deduced that hc for as-depositedlms corresponds to both the spherical and conical regions of theerkovich indenter. This argument is furthermore supported by

926 A. Chatterjee et al. / Wear 265 (2008) 921–929

) as-d

twoftttiov

wfitFc

Fig. 5. AFM surface morphologies of residual scratches of (a

he fact that the present hardness value (22 GPa) falls inside theindow of the hardness values reported earlier (21–25 GPa) [20]n similar but thicker films (∼500 nm). On the other hand, hcor annealed films (8 nm) is comparable to the spherical–conicalransition depth (8.78 nm). Therefore, it is reasonable to expecthat deformation is imposed by the spherical region of the inden-

er. At such low contact depths, mechanical response of the films expected to be more elastic than plastic. A higher componentf elastic response may account for the slightly higher hardnessalue (43 GPa) compared to the one reported earlier (40 GPa),

rape

eposited HfB2 film, (b) annealed HfB2 film and (c) TiN film.

here indentations were performed up to depths ∼50 nm. Thelms in our case are thinner and, to conform to the 10% inden-

ation rule, our indentation depths had therefore to be shallower.or better accuracy at such shallow depths, one could use aube-corner tip.

From nanoscratch data, we also measured a change in the

elative elastic–plastic response of the films before and afternnealing. Hertzian contact pressure calculations and their com-arison with yield strengths show that plastic deformation isxpected for as-deposited HfB2 and TiN films while elastic

A. Chatterjee et al. / Wear 265 (2008) 921–929 927

Fig. 6. Variation of μ with normal load at sliding velocity 0.32 �m/s, for 1 pass.Here and in the following, error bars represent one standard deviation obtainedfrom 6 to 10 independent measurements.

Fig. 7. AFM morphology of residual scratches in TiN when the sliding velocitywas varied for a normal load of 200 �N.

Fig. 8. Variation of μ with sliding velocity at a normal load of 500 �N, for 1pass.

Fig. 9. Variation of μ with number of passes at normal load 500 �N, slidingvelocity 0.32 �m/s.

Fig. 10. Variation of contact depths and contact areas with number of passes atnormal load 500 �N, sliding velocity 0.32 �m/s.

Fig. 11. Variation of percentage change in contact area (�A) and percentagechange in friction coefficient (�μ) vs. number of passes at normal load 500 �N,sliding velocity 0.32 �m/s.

9 Wear

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28 A. Chatterjee et al. /

esponse is excepted of annealed HfB2, at least for the 100nd 200 �N normal load experiments. Preliminary finite elementnalysis (FEA) studies of nanoindentation of hafnium diboridelms indicate that yield strength values corresponding to theest-fit simulation curve of experimental load–displacementata satisfy H/Y = 2.1. While this might seem to be a deviationrom the well-known Tabor’s equation, H = 2.7Y it is consistentith the linear dependence of H/Y versus E/Y as reported byabor [27]. A ratio of H/Y = 3 is valid only for materials with/Y ∼ 100. For hafnium diboride, E/Y ∼ 24.7 and corresponding

o this ratio, H/Y = 2.1 is indeed expected from Tabor’s work./Y = 2.1 has been reported for TiN films [28] whose E/Y val-es were similar to our films. If one uses the relation H/Y = 2.1or obtaining yield strengths from hardness calculations, theertzian contact pressures of annealed HfB2 will be lower than

he yield strength for the entire range of experiments (from 100o 500 �N) and therefore elastic response is expected at all loads.urthermore, elastic response would also be expected of as-eposited hafnium diboride at 100 �N. In summary, while theres some uncertainty on the exact yield strength values, the mainonclusions from the Hertzian pressure analysis remain that,or the loads investigated here, the response of the as-depositedaterial is largely plastic, while that of the annealed film should

ave a dominant elastic component.A different analysis on the same set of nanoscratch results

ased on the phenomenological model proposed by Moscht al. [13] leads to similar conclusions. It predicts the rela-ive elastic–plastic response based on the Mayer’s constant (n)xtraction; n values of 1.78 for TiN and 1.71 for as-depositedfB2 indicate that for these two films, plastic deformation domi-ated but for annealed films, n = 2.43 which indicates a dominantlastic response. The resistance of annealed films to plowing cor-esponds well with their higher hardness values obtained fromanoindentation experiments as well as a previous study of theechanical properties of as-deposited and annealed HfB2 films

20]. Hardness is essentially a measure of resistance to plasticeformation and so harder materials will have inherently higheresistance to plowing. Thus the friction behavior of the films isonsistent with their mechanical properties.

The variation of friction coefficient and contact area withhe number of passes reveals interesting differences betweens-deposited and annealed HfB2. Friction force is defined by= σshearAcontact, where σshear is the shear stress and Acontact is

he contact area. An increase in the number of passes might leado a progressively increasing wear of the material leading toigher penetration of the tip in the material. In such a case, theontact depth and contact area increase progressively with theumber of passes. As shown in Fig. 10, as-deposited HfB2 showshis effect. The similarity in the percentage change in contactrea (�A) versus number of passes to the percentage change inriction coefficient (�μ) for as-deposited films strongly suggestshat increase in the friction coefficient with the number of passess primarily a manifestation of the increase of the contact area

or as-deposited HfB2. This explanation is also consistent withhe fact that as-deposited HfB2 is expected to deform plasticallynd therefore less likely to recover the change in the deformationone with increasing number of passes.

•

265 (2008) 921–929

While the “change in contact area” approach does indeedxplain the variation in friction coefficient for as-depositedfB2, it does not explain the friction variation with the num-er of passes for annealed HfB2 for which �μ (87.5% at 50asses) is much greater than �A (11.5% at 50 passes). Moreover,significant component of the mechanical response of annealedfB2 to deformation is expected to be elastic. Thus we shouldot expect an appreciable change in �A with increasing numberf passes. The increase in �μ for annealed HfB2 however cane explained by the running-in behavior of the film–tip interface29]. During the initial stages of running-in, due to interactionetween the diamond tip and the opposing asperities on the film,he friction force rises. It stabilizes only after the asperities ofhe film break down. For annealed HfB2 films, it is apparent thatven up to 50 passes, the friction force has not stabilized and soll the asperities have not yet been run-in.

Achanta et al. [16] have reported that the debris created due tosperity-interaction dominates over the rolling effect of ejectedear particles and is responsible for an increase in the coeffi-

ient of friction with the number of passes, for both DLC andiN coatings. Since similar trends are observed in our experi-ents, we propose that the increase in the friction coefficient

f annealed HfB2 and the prolonged presence of the running-ineriod is due to the diminishing effect of the so-called wear rolls.he rolling effect due to the ejected particles can be rational-

zed by a sink–source model [17]. At higher number of passes,he surface smoothens which results in a decline in the rate ofarticle formation (source). On the other hand, high contact pres-ures ensure that the fragmentation of existing particles becomesore pronounced as the number of passes increase. These small

ragments get trapped in the inter-asperity region of the coatingsink). Thus, the relative effect of source gets diminished andhat of sink gets enhanced, resulting in an increase in the fric-ion coefficient of annealed films with increase with the numberf sliding passes.

. Conclusions

We have explored the nanofriction behavior of CVD-grownafnium diboride films, in as-deposited and annealed condi-ions, for normal loads ranging from 100 to 500 �N. The mainonclusions that can be drawn are:

As-deposited and annealed films show favorable fric-tion behavior, with a lower average friction coefficient(0.097 ± 0.001 for as-deposited HfB2 and 0.087 ± 0.002 forannealed HfB2), compared to TiN (0.156 ± 0.006).The dependence of the friction coefficient with normal load,supplemented by Meyer’s constant calculations, suggest thatas-deposited HfB2 and TiN show uniformly plastic responseat all normal loads used in the experiments. Annealed HfB2exhibit elastic behavior which is supported by the comparisonof Hertzian contact pressure and yield strength.

The plowing mechanism dominates the friction behavior ofas-deposited HfB2 and TiN. For annealed films, however,the negative slope of the friction versus load curve indicatesclearly that the adhesive mechanism dominates. This is sup-

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ported by the near-constant nanoscratch groove depths forloads less than 500 �N.None of the films show any significant dependence on slid-ing velocity, thus indicating that flash temperature heating ofasperities is negligible.Increase in friction coefficient in as-deposited HfB2 withincrease in the number of passes is a result of accompanyingincrease in contact areas.Increase in friction coefficients in annealed HfB2 is a result ofthe decreasing effect of wear-rolls and indicates that all filmasperities have not been run-in even after 50 passes.

The present results suggest that the HfB2 thin films investi-ated here, in particular annealed films, should have very goodear resistance at the nanoscale. AFM-based nanoscale wear

xperiments are currently being performed to confirm this pre-iction.

cknowledgements

The authors are grateful to the National Science Founda-ion and CNRS, France, for support of this research under grantumber NSF DMR-0354060 through a NSF-CNRS cooperativegreement. Compositional and structural analyses of the filmsere carried out in the Center for Microanalysis of Materials,niversity of Illinois, which is partially supported by the U.S.epartment of Energy under grant DEFG02-91-ER45439. We

hank Prof. J.-P. Chevalier for numerous stimulating discussions.e thank Profs. Joe Greene and Ivan Petrov for providing us with

he TiN coating used as a point of comparison in this work.

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