Int. Journal of Refractory Metals & Hard Materials 28 (2010) 15–22

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Int. Journal of Refractory Metals & Hard Materials

journal homepage: www.elsevier .com/locate / IJRMHM

Structure, mechanical properties and tribology of W–N and W–O coatings

T. Polcar a,*, A. Cavaleiro b

a Department of Control Engineering, Faculty of Electrical Engineering, Czech Technical University in Prague, Technická 2, Prague 6, Czech Republicb SEG-CEMUC – Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, P-3030 788 Coimbra, Portugal

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

Article history:Received 21 February 2009Accepted 13 July 2009

Keywords:High temperature tribologyTungsten nitrideTungsten oxideProtective coatings

0263-4368/$ - see front matter � 2009 Elsevier Ltd. Adoi:10.1016/j.ijrmhm.2009.07.013

* Corresponding author. Tel./fax: +420 224 312 439E-mail address: polcar@fel.cvut.cz (T. Polcar).

The tribological properties of nitrides and oxides of transition metal thin films deposited by reactive mag-netron sputtering have been thoroughly studied for three decades. Nevertheless, there are still severalgaps in knowledge. The majority of studies are focused on a limited number of metals, namely Ti, Aland Cr, while other potentially attractive compounds are aside the main attention. Even in case of TiN,probably the most studied hard thin film, the frictional and wear behaviour brings many controversies.Despite significant progress of analytical and computational methods, the analysis of the wear behaviouris still a great challenge.

We are presenting here a summary of our recent work on tungsten nitride (nitrogen content 0–58 at.%)and tungsten oxide (oxygen content 0–75 at.%) coatings deposited by reactive magnetron sputtering. Ouraim has been the analysis of the connection of fundamental properties of these films, such as chemicalcomposition, structure, hardness, Young’s modulus and residual stress, with their tribological properties– friction coefficient and wear rate. We have been focused mainly on the description of the dominantwear mechanisms influencing the tribological properties. The tribological tests have been carried outboth at room and elevated temperature; the temperature was increased in steps until immediate coatingsfailure.

The tungsten nitride coatings with the ‘‘worst” parameters generally considered as vital for high wearresistance, such as hardness, were considered to have the best tribological performance.

� 2009 Elsevier Ltd. All rights reserved.

1. Introduction

Coatings of transition metals nitrides are known for their highhardness and excellent wear resistance, being titanium nitridethe best known member of this family. However, its tribologicalproperties are significantly diminished when the coating is ex-posed to elevated temperatures, particularly due to rapid oxida-tion. In this respect, chromium nitride possesses significantadvantages compared to TiN, mainly higher oxidation resistance,thermal stability and, particularly, corrosion resistance. On theother hand, as main drawback CrN exhibits a significantly higherfriction coefficient, which is not desired in many applications [1].Other nitrides, such as MoN, VN or AlN, could outperform TiN orCrN in many parameters; however, their commercial use is verylimited. Tungsten nitride stays still aside of attention despite somesignificant advantages, such as extremely high hardness and excel-lent adhesion on typical steel substrates.

ll rights reserved.

.

W–N coatings were studied for mechanical applications [2];nevertheless, the main attention was paid to ternary systems pre-pared by the addition of different elements (e.g., W–Ti–N [3] andW–Si–N [4]), with the final scope of improving the mechanicalproperties of the coatings. Tribological properties of tungsten ni-tride coatings at room [2,5] and high [6] temperature were thor-oughly studied by present authors. It has been shown that thetribological properties of tungsten nitride coatings are stronglyinfluenced by the formation of a tungsten oxide tribolayer evenat room temperature. Further increase of environment tempera-ture led to the oxidation of the coatings, particularly in the areaof the contact. A great amount of research work was also carriedout by the authors in the W–O system [7,8]. Besides a completeand detailed basic characterization of the coatings, particular inter-est was paid to their thermal stability in both protective andoxidant atmospheres [9]. In some specific cases, the overall tribo-logical behaviour of the coatings was assessed by pin-on-diskagainst different counterbody materials [5–7].

This paper summarizes our work dealing with the tribologicalproperties of tungsten nitrides deposited by magnetron sputtering.Additional results concerned with the W–O system will be alsopresented as a tool for interpreting and understanding the globaltribological behaviour of these systems.

10 30 50 70

Inte

nsity

(arb

. uni

ts)

2 Theta (deg)

W100W90N10

W88N12

W45N55

W81O13

W25O75 *

*

*

++

+

ºº

ºº

ºº

º

#

&

#

&

Fig. 1a. Typical X-ray diffraction patterns of W–N and W–O sputtered coatings.(* – Substrate; + – WO3; o – b-W, # – W2N, and – a-W).

Fig. 1b. HR-TEM image of W26O74 coating.

16 T. Polcar, A. Cavaleiro / Int. Journal of Refractory Metals & Hard Materials 28 (2010) 15–22

2. Experimental details

Tungsten oxide and nitride coatings were deposited by DC reac-tive magnetron sputtering from a tungsten target in a Ar + O2 or N2

atmosphere, respectively, onto high speed steel AISI M2 substratesheat treated to have a hardness close to 9 GPa. The deposition runsfor both tungsten compounds were performed keeping constantthe following parameters: total working pressure (0.3 Pa), targetcurrent density (10 mA cm�2), substrate at floating potential, nosubstrate rotation, inter-electrode distance (65 mm) and substratetemperature (<350 �C, with no external heating). In order toachieve different N and O contents in the films, the coatings weredeposited with different flow ratios of the reactive gas, N2/Ar from0 to 3 and O2/Ar from 0 to 1, respectively. Before any deposition, anultimate vacuum pressure better than 5 � 10�4 Pa was reachedand the substrates surface was ion cleaned with an ion gun.

The chemical composition and the structure were determinedby a Cameca SX-50 electron probe microanalysis apparatus (EPMA)and an X-Pert Philips X-ray diffractometer (XRD), respectively. Theoxides microstructure was analyzed by high resolution transmis-sion electron microscope (HR-TEM – JEOL 2010-FEG, operated at200 kV). The mechanical properties of the coatings were evaluatedby depth-sensing indentation technique using a Fischer Instru-ments-Fischerscope, with a maximum load of 50 mN and followingthe testing procedure described elsewhere [10]. The adhesion/cohesion of the coatings was evaluated by scratch-testing tech-nique using a Revetest, CSM Instruments. The load was increasedlinearly from 0 to 50 N, using a Rockwell C 200 lm radius indentertip, loading rate of 100 N/min, and scratching speed of 10 mm/min.

Wear testing was done using a high temperature pin-on-disctribometer (CSEM Instruments). Al2O3 balls with a diameter of6 mm were used as counter-parts. All measurements were pro-vided with a load of 5 N and a linear speed of 0.05 m/s; the relativehumidity of the air was kept constant at (25 ± 5)%. The maximumstatic Hertzian pressure for the elastic contact between the balland the coating was about 1.5 GPa. The worn volume of the ballswas evaluated by optical microscopy. The morphology of the coat-ing surface, ball scars, wear tracks and wear debris were examinedby scanning electron microscopy (SEM); the chemical analysis ofthe wear tracks and the wear debris was obtained by energy-dis-persive X-ray analysis (EDX). The profiles of the wear tracks weremeasured by mechanical profilometer. The wear rates of the balland the coating were calculated according to [6] as the worn mate-rial volume per sliding distance and normal load. The average valueof five profiles measured on each wear track was used to calculatethe coating wear rate.

3. Results and discussion

3.1. Coating characterization at room temperature

3.1.1. Chemical composition and structureThe increase of the nitrogen partial pressure ratio ðpN2

=pArÞ from0 to 3 led to a linear increase of the nitrogen content in the filmsfrom 0 to 55 at.%. The oxygen content originating from the residualatmosphere and the chamber leaks was lower than 5 at.%; whenhigher nitrogen flows were used, the oxygen content decreased.Similar results were obtained for the W–O system, but being pos-sible to reach much higher O contents (�75 at.%) even if the max-imum used partial pressure ratio was much lower ðpO2

=pAr ¼ 1Þ.This is due to the much higher affinity of W for O than for N. Itshould be pointed out that the low reactivity of tungsten allowssputtering of nitrides and oxides in all ranges of chemicalcompositions.

Typical XRD diffractograms for W–N and W–O are shown inFig. 1a [5,8]. The structure of the W–N and W–O coatings can becorrelated with their chemical composition: (i) for low-nitrogenor oxygen contents (N,O < 9 at.%), the coatings exhibited the equi-librium bcc a-W phase with a (1 1 0) preferential orientation; (ii)for coatings with intermediate nitrogen and oxygen contents upto 30 at.% (e.g., W88N12, W85N15 or W81O19) the b-W phase is de-tected mixed with the a-W; in the N-containing films for at.%N > 15at.% fcc. NaCl-type b-W2N is also occurring (iii) for coatingswith higher reactive element contents (N,O > 30 at.%) the structureis either formed by the NaCl-type nitride compound, b-W2N, with(2 2 0) preferential orientation, in W–N coatings or by a low crys-tallinity order phase. It should be pointed out that an increase ofnitrogen content over 55 at.% was followed by the formation ofhexagonal d-WN phase [6] whereas the structure of the W–O coat-ings with high oxygen content can be described as nanocrystallineor quasi-amorphous. Fig. 1b shows a high resolution TEM image ofthe microstructure of this film. The presence of very small crystal-lites embedded in a low order crystallinity matrix is clearly obser-vable. SEM micrographs of all W–N coating cross-sections showedtypical columnar morphology contrasting with the featurelessW–O films (Fig. 2).

Fig. 2. SEM cross-sections of W65N35 (a) and W65O35 (b) coatings (Si substrates).

0 10 20 30 40 50 60 700

4

8

12

16

20

24

28

32

36

40

44 W-N - Hardness W-O - Hardness W-N - Res. stress W-O - Res. stress

N or O content (at.%)

Har

dnes

s (G

Pa)

-7

-6

-5

-4

-3

-2

-1

0

1

2

3

4

5

Res

idua

l str

ess

(GPa

)

Fig. 3. Hardness and residual stress of W–N and W–O coatings with differentnitrogen and oxygen content, respectively.

0 10 20 30 40 50 60 70 80

0

5

10

15

20

25

30

35

40

45

W-N W-O

Crit

ical

load

Lc2

(N)

N or O content (at.%)

Fig. 4. Critical load Lc2 for W–N and W–O coatings.

T. Polcar, A. Cavaleiro / Int. Journal of Refractory Metals & Hard Materials 28 (2010) 15–22 17

3.1.2. Hardness, residual stress, and adhesionThe evolution of the residual stress and the hardness as a func-

tion of the nitrogen content in W–N coatings, depicted in Fig. 3, fol-lows similar trends. A linear relationship can be establishedbetween the hardness and the residual stress, showing that the lat-tice distortions related to the structural defects that induceincreasing intrinsic stresses are determinant effects influencingthe measured hardness of the coating.

All coatings display a compressive residual stress state, result-ing from the backscattered energetic neutral impinging on thegrowing film surface [11]. The increasing nitrogen content from 0to 9 at.% promotes a shift in the a-W (1 1 0) peak position to lower

diffraction angles. Therefore, the nitrogen is probably being placedin interstitial positions in the lattice increasing the dilatation of thelattice parameter and inducing increasing compressive stressesfrom 3.7 to 4.8 GPa. Consequently, the hardness also increasesfrom 22 to 40 GPa. In the case of W–O coatings the effect of theinterstitial position of oxygen in the W lattice does not give riseto any hardening effect. Firstly, the distortion caused in the W lat-tice is much lower for O than for N (e.g., for �10 at.% N(O), 0.9% forO against 2.5% of N [8,12]) and, secondly, the higher electronega-tivity of O in relation to N decreases the covalent character of thebond, giving rise to lower materials strength. b-W phase, whichis dominant for W88N12 and W84N16 coatings, is known to havelow compressive residual stress [13], or even tensile stress [14].The residual stress value of these two coatings dropped to �3.3and �3.4 GPa, respectively, leading to a decrease in the hardnessdown to values close to 29 GPa. On the other hand, the occurrenceof this phase in the W–O coating gives rise to a small increase inthe hardness (W87O13 film).

For high nitrogen content (N P 34 at.%), the coatings display theb-W2N phase, without or with (2 2 0) preferential orientation. Thehardness of these coatings decreases proportionally from 41 to30 GPa. It can be interpreted by the synergetic effect of the com-pressive residual stress decrease with the change in the preferen-tial orientation of the films. In the W–O films the hardnesslinearly decreases from 25 (W87O13) down to 7.7 GPa (W25O75).The amorphisation of the coatings associated with the increasingimportance of the W–O ionic bonding type justifies the drop inthe coatings hardness.

The adhesion/cohesion values of the coatings were measured byscratch-test apparatus to calculate the critical loads Lc1 and Lc2

(see Fig. 4). These critical loads are defined as the first cohesiveand adhesive failures, respectively. With increasing nitrogen con-tent Lc1 and Lc2 values decreased linearly from 18 to 9 N and from50 to 13 N, respectively. Pure tungsten coating exhibited Lc2 valuehigher than the maximum applied load (50 N). The scratch-estbehaviour is even worse for W–O films, Lc2 decreases abruptlydown to �10 N for O content up to 30 at.% and is kept approxi-mately constant for higher contents (W25O75 film �9 N).

3.1.3. Friction and wear of the low-Nitrogen W–N coatingsFig. 5 shows the evolution of the friction coefficient and the

coating wear rate with increasing nitrogen content when the coat-ings are sliding against Al2O3 balls. The main feature which seemsto determine the inversion in the tribological behaviour of W–Ncoatings is the presence of the NaCl-type W2N phase. In fact, thedrop in the friction and wear coefficients was registered for the

0 10 20 30 40 50 600.2

0.3

0.4

0.5

0.6

0.7

Nitrogen content (at.%)

Fric

tion

coef

ficie

nt

0.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

4.0β-W2Nδ-WN

β-W2Nα-W + β-W2N

Wea

r rat

e (1

0-6 m

m3 /N

m)

α-W

Fig. 5. Tribological properties of W–N films at room temperature. The dominantstructure zones are indicated.

Fig. 6. Typical wear debris taken from the ball, tests with low-nitrogen coatings.The large particles are clearly visible.

Fig. 7. SEM micrograph of the wear track, coating W88N12 (back-scattered mode).The defects are clearly visible (see black arrows).

Fig. 8. SEM micrograph of W85N15 wear track.

18 T. Polcar, A. Cavaleiro / Int. Journal of Refractory Metals & Hard Materials 28 (2010) 15–22

coating in which this phase started to be dominant. Al2O3 ball scarswere very small and transfer of the coating material to the ball sur-face has not been observed abundantly; the only exception was thetest with W88N12 film, where a measurable ball scar was detected.

SEM analysis of the wear particles has revealed that there aretwo different types of wear particles: small round particles withan average dimension of less than 50 nm and large particles withsharp edges with a grit size of 1–3 lm (Fig. 6). The highest concen-tration of the large particles occurs in the wear debris of the coat-ing with 12 at.% N; however, even in this case, the amount of thesmall particles is clearly dominant. The chemical composition ofthe small particles measured by EDX is almost stoichiometric tung-sten trioxide with no vestiges of nitrogen, while the large particlesare not oxidized.

The wear track of W95N5 coating shows only shallow scratchesparallel to the relative sliding movement. Defects in the weartracks could be observed by SEM in the films with further increaseof nitrogen content. The W90N10 wear track is very similar to theprevious one; however, cohesive failures of the coating could bedetected. These failures contributed to the increase of the wearrate, since the wear debris containing large tungsten nitride parti-cles can cause abrasive wear due to their sharp edges and highhardness. The W88N12 sample exhibits the highest wear rate andfriction coefficient. The number of delaminations in the wear trackrapidly increases as well as the number of large W–N particles,causing very deep scratches in the coating surface. The largestscratches are finished by the accumulation of coating materialrevealing plastic deformation (Fig. 7). It can be assumed that alarge particle may get jammed in the coating surface causing a lo-

cal increase of temperature, which enhances the tribo-oxidation ofthe coating. The tungsten oxides are much softer than the nitridesand, thus, local plastic deformation can occur. Oxidation of mate-rial in the defect was confirmed by EDX. This coating exhibits aswell the lowest value of the critical load Lc1, which can enhancethe formation of large delaminated particles. Moreover, the ballwear scar is covered with deep scratches and, thus, the delaminat-ed particles are abrasive enough even to damage the hard Al2O3

ball surface. No major defects of plastic deformations are visiblein the wear track of the W85N15 coating (Fig. 8). As a consequence,compared to W88N12 sample, the wear rate and the friction coeffi-cient decrease significantly. This different behaviour can not beattributed to any change in the mechanical properties, since hard-ness, residual stress and critical load of both coatings are almostidentical. However, as referred to above, the XRD analysis has re-vealed that from one coating to the other, the presence of the b-W2N phase was detected which seems to affect the tribologicalbehaviour. In fact, the coatings formed predominantly by the b-W2N phase present similar wear mechanisms as those of W85N15

one.In general, for low-nitrogen content films (N content < 15 at.%),

the wear is driven by a combination of abrasion, delamination andtribo-oxidation. The influence of the interlayer of wear debris,which is formed between the surfaces in the contact zone, is neg-ligible, since it does not adhere to any surface and is rapidly drivenout from the contact area. Tribo-oxidation occurs during the entiretest, since the majority of the wear debris particles are oxidized.The frictional heat generated by dry sliding is almost exclusivelydissipated away through the asperities in contact causing a signif-

T. Polcar, A. Cavaleiro / Int. Journal of Refractory Metals & Hard Materials 28 (2010) 15–22 19

icant local increase of temperature leading to the oxidation of eachasperity [15]. Thus, during the contact, as tungsten oxides havelower density than the nitrides, the asperities will oxidize and ex-pand lifting apart the remaining asperities. This effect results in theconcentration of both the frictional energy dissipation and themechanical load on a few asperities only. The tungsten oxidesare soft and the oxidized asperities can be plastically deformedand easily removed from the coating surface. This mechanism istypical for W95N5 coating and can be described as a polishingand mild wear mode.

The delamination of large particles leads to tungsten-rich (i.e.,not oxidized) particles in the wear debris for W90N10 coatingsrevealing that other mechanism of material removal has to be con-sidered. This coating exhibits the highest hardness, Young’s modu-lus and, particularly, residual stress values. It is possible that theexpansion of the oxidized asperities can not be accommodatedby the coating and, together with the stresses induced by the fric-tion, creates cracks and enhances their propagation leading tocoating chipping. The presence of hard sharp-edged particles inthe contact is typical of abrasive wear, as it is confirmed by thepresence of scratches in the wear track. The wear and friction coef-ficient increases as a result of the presence of the non-oxidized W–N particles in the contact. Therefore, the abrasive process anddelamination are the main wear mechanisms for W90N10 film.These mechanisms are complemented by local plastic deformationin the case of W88N12. In this case, in spite of having a global lowerstress state due to the presence of b-W phase, the delamination canstill occur due to either the high local stresses in the a-W phase(see Fig. 3) or due to dimensional changes in b-W during sliding.As the hardness is lower, the large particles can be embedded inthe coating surface and be plastically deformed.

W85N15 coating stands for the transition in the wear mechanismfrom low-N to high-N content films. In fact, in spite of havingapproximately the same mechanical properties (H, E, Lc, rr) asW88N12 film, its tribological behaviour starts to be different. Com-pared to W88N12 films, the structure of the W85N15 coating is differ-ent: (i) the amount of b-W phase decreases, (ii) the a-W phaserecovers, with a significant decrease in its lattice distortion and lo-cal stresses, and (iii) W2N phase appears. The combine effect ofthese structural changes results in similar mechanical but differenttribological properties. Neither plastic deformation defects in thewear track nor W–N particles in the wear debris were detectedafter the running test. The aspect of the wear track is only of pol-ishing wear, resulting in a significant decrease of the friction coef-ficient and the wear rate. Moreover, for this film no adhered platescould be observed in the wear track.

3.1.4. Sliding of W–N coatings with high nitrogen contentThe wear debris produced during sliding of high-N content films

is very homogenous containing only small round particles with adiameter of 20–30 nm. Oxidation studies performed in W–N coat-ings showed that, in coatings with high N content, the oxide scalespalled off as it was being formed whereas it stayed attached to thebase coating in the case of low-N coatings. This can explain theeasy formation of the very small wear particles in these coatingscompared to the low-N content ones. Microanalysis of the weardebris has shown that their chemical composition is stoichiometrictungsten trioxide. The volume of adherent wear debris is muchhigher compared to that of low N-content coatings. The weartracks are very smooth, neither defects nor cracks have been ob-served. The large volume of adhered wear debris, which consider-ably exceeds the worn volume of the coating, can be easilyexplained by the difference in the density of the coatings and theadhered wear debris. The wear debris is exclusively stoichiometrictungsten trioxide, which is much less dense than bulk WO3 with adensity of only 7.16 g cm�3 [16], when compared to 16.2 g cm�3

for the W–Nitride calculated from XRD patterns. In case ofW45N55 the wear debris covers almost all the wear track.

The sliding of the coatings with high nitrogen content is typicalof a three-body abrasive wear. In this case, the wear particles arefree to slide over the surface, since they are not held rigidly. Con-sequently, the wear rate is much lower than for a two-body wear,characteristic of the low-N content films. The three-body wear isallowed by a rapid production of the wear debris in the first stagesof the sliding. The thick interlayer between opposing surfaces isresponsible for the increase of the contact area resulting in a de-crease of the contact pressure. Consequently, the wear rate andthe friction coefficient decrease. Moreover, the wear debris adhereson the sides of the wear track, where the contact pressure reachesminimal values. The surface of the central groove of the wear track,which is neither covered by adhered oxide layer nor oxidized, isthe only area of the coating material exposed to direct wear, sinceadhered oxide layers totally protect the coating surface.

3.1.5. Sliding of selected W–O coatingsWhen the tribological behaviour of the W–O coatings is ana-

lyzed and compared to W–N films it can be immediately concludedthat their performance is inferior concerning both the friction coef-ficient and the wear rates. For example, for low interstitial contentfilms W88N12 and W87O13 the values are l = 0.6 and 0.73 and thewear rate 3.3 and 4.4 � 10�6 mm3/Nm, respectively, whereas forcompound films W45N55 and W25O75 the values are l = 0.33 and0.49 and the wear rate 0.15 and 2.9 � 10�6 mm3/Nm, respectively[5,7]. In the low O content film, also containing the b-W phase, inspite of its moderate hardness (�25 GPa) no formation of big wearparticles takes place due to the low local stresses in the a-W phase.Thus, the wear track is rather smooth with only shallow scratches.However, local well adherent plates of tungsten oxide cover par-tially the wear track. For the W25O75 coating a high amount of weardebris is formed promoting a very compact and homogeneouslayer covering entirely the wear track. Moreover, also on the ballsurface tungsten oxide can be detected. In this case, an almostexclusive third-body wear mechanism prevails with self-matedtungsten oxide.

At a first sight it seems that in both W–N and W–O systems thesliding mechanisms are quite similar, i.e., for low interstitial con-tents the contact between the ball and the coating is direct, with-out or with low amount of third-body in-between, whereas forhigh N/O contents a high amount of wear debris are present inthe sliding contact. Therefore, the differences in the friction coeffi-cients for both systems are surprising, particularly when the fric-tion of oxides is higher than that of nitrides – quite unusualunder similar conditions. Particularly the case with low-N/O con-tent is difficult to explain, since the typical abrasive wear forW88N12 coating should exhibit higher friction. The possible reasonfor higher friction of W87O13 can be the adhesion of the tungstenoxide tribolayer to the transferred oxidized coating material onthe alumina ball. Similar process could take place in sliding ofW25O75. Alumina balls sliding against nitrides, on the other hand,showed not vestiges of adhered coating material. In other words,the steady-state sliding of W45N55 film is represented mainly byAl2O3 ball – W–O tribolayer contact, whereas self-mating of tung-sten oxides become dominant after running in with W25O75 coat-ing, which can lead to strong adhesive forces.

3.2. Characterization of WN films at elevated temperature

3.2.1. Structure and hardnessAs referred to above, the tribological behaviour of W–N coatings

at room temperature could be divided into two groups with differ-ent dominant wear mechanisms. The coatings with low-nitrogencontent (up to 15 at.%) exhibited higher values of friction coeffi-

0 100 200 300 400 500 6000

5

10

15

20

25

30

35

40

45

Har

dnes

s (G

Pa)

Annealing temperature (ºC)

W70

N30

W53

N47

W42

N58

a

200

250

300

350

400

450

You

ng's

mod

ulus

(G

Pa)

W30

N70

W53

N47

b

20 T. Polcar, A. Cavaleiro / Int. Journal of Refractory Metals & Hard Materials 28 (2010) 15–22

cient and wear rate, since their sliding properties were dominatedby abrasive wear and delamination. High nitrogen contents (33–55 at.%) led to a typical three-body wear with a thick layer of tung-sten trioxide consisting of small round particles. Considering thehardness and the wear rate of tungsten nitrides, three chemicalcompositions from the later group were selected for the tribologi-cal tests at elevated temperatures due to the following reasons: (i)the coatings with high nitrogen content exhibited lower wear, (ii)the high wear resistance of these coatings was caused by the for-mation of a third-body of tungsten trioxide; this phenomenonshould be enhanced at higher temperatures, and (iii) the selectedchemical compositions cover coatings with either the highesthardness or the lowest friction and wear.

All coatings showed similar diffraction patterns from room tem-perature up to 400 �C, see Fig. 9. Only a small shift of the XRD peakspositions to higher diffraction angles can be observed being attrib-uted to stress relaxation. The presence of small oxide peaks close to2h � 25� revealed that at 500 �C surface oxidation already occurredwhile, at 600 �C, XRD patterns of the coatings after annealingexhibited only the m-WO3 structure (WO3 monoclinic phase, ICDDcard nr. 83-0950) with (0 0 2) preferential orientation. The absenceof any signs of nitride phases clearly demonstrated the entire oxi-dation of the coating. It should be pointed out that the XRD pat-terns obtained after annealing at 600 �C were identical for all thetested coatings.

Finally, the coating series deposited with the highest nitrogencontent, W42N58, clearly exhibited at room temperature a mixtureof d-WN and b-W2N phases (it is not clear from XRD spectra ofW45N55 coating [5] whether d-WN is present or not).

The variation of the hardness and the Young’s modulus with theannealing temperature is shown in Fig. 10. At room temperature,

150

W42

N58

0 100 200 300 400 500 600

Annealing temperature (ºC)

Fig. 10. Hardness (a) and Young’s modulus (b) of annealed W–N coatings.

Fig. 9. XRD of W–N coatings selected for high temperature tribology tests.

W70N30 was the hardest coating (�41 GPa), while the coatings withhigher nitrogen content, W53N47 and W42N58, exhibited similarhardness values close to 30 GPa (see more details above). Withincreasing annealing temperature up to 400 �C a slow decrease ofthe hardness could be observed. The difference in H valuesbetween the room temperature and 400 �C was approximately8 GPa in all the three coatings. Since neither variation of the struc-ture nor oxidation was observed, the hardness decrease with tem-perature can be attributed to a decrease of the compressive stress,which should be closely related to the decrease of defect densityinduced by the heating and the consequent stress relaxation. Theseeffects complemented with the start of the oxidation process jus-tify the steeper drop in the hardness observed for higher tempera-tures. As shown in previous section, the coatings are fully oxidizedat 600 �C. As a consequence, at this temperature the hardness isconstant for all the coatings, 6 GPa, value typical for the tungstentrioxide [7].

The coatings W70N30 and W53N47 exhibited similar E values upto 400 �C, close to 380 GPa. At 500 �C, the Young’s modulus slightlydecreased to 310 and 340 GPa, respectively, and at 600 �C droppeddown to 180 GPa, again a value corresponding to WO3 [8]. TheYoung’s modulus of the coating with the highest nitrogen content,W42N58, was lower than that of the other coatings.

3.2.2. Friction and wearThe average friction coefficient and the wear rate as a function

of the testing temperature of the coatings are presented in Fig. 11.Both the friction and the wear exhibited similar trends for all the

0 100 200 300 400 5000.0

0.2

0.4

0.6

0.8

1.0 W

70N

30 W

53N

47 W

42N

58

W70

N30

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N47

W42

N58

Temperature [ oC]

Fri

ctio

n c

oef

fici

ent

0

1

2

3

4

5

6

7

Wea

r ra

te [

10-6m

m3N

-1m

-1]

Fig. 11. Friction and wear rate of W–N coatings as a function of the testingtemperature. Open symbols represent friction coefficient, full ones wear rate.

Fig. 13. Cracks in the wear track, W70N30 coating, tested at temperature 500 �C.

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coatings. The friction increased from the lowest level at room tem-perature to a maximum in the temperature range from 200 to300 �C. A further rise in the temperature was followed by a smalldecrease of the friction. Considering the error of the measurement,the friction value was almost identical at 300 and 400 �C for all thenitrogen contents, while the lowest friction was exhibited by theW42N58 coating at temperatures up to 200 �C and by W70N30 coat-ing at the temperature of 500 �C. The worn volume was almost notmeasurable up to 200 �C and then increased for higher tempera-tures. All the coatings tested at 600 �C were worn out after the test;thus, the friction and the wear data are not presented.

As referred to above, the wear tracks produced at room temper-ature showed a thick layer of tungsten oxide adhered on their bor-ders; only on the central groove, the direct contact between the balland the tungsten nitride took place. The adhered layer on the sidestarted to disappear with increasing temperature and vanished at300 �C. The central part of the wear tracks, uncovered at room tem-perature, became covered by isolated isles of tungsten oxide at200 �C (Fig. 12). The wear tracks produced on the different tungstennitrides were very similar in the range of 20–200 �C, while signifi-cant differences were observed at higher temperatures.

The wear track produced on the coating W70N30 when sliding at300 �C exhibits small lateral cracks originated from the centre of

Fig. 12. Wear track of W70N30 coating produced by sliding at 200 �C (SEM). Note theadhered tribolayer in the centre of the track. Other two tested WN coatingsexhibited similar appearance at this temperature.

the wear track. The cracks covered the entire wear track at500 �C (Fig. 13), and the iron signal detected by EDX revealed thelocal destruction of the coating. Therefore, the wear rate for500 �C is higher than that indicated in Fig. 11. The wear tracks ofW53N47 and W42N58 coatings were almost identical in all the tem-perature range. At 300 and 400 �C, the first damages appeared inthe wear tracks. The partial destruction of the coating always oc-curred under the adhered tungsten oxide layer. A further rise inthe temperature to 500 �C brought a radical change, since theadhesive tungsten oxide vanished and long straps showing a se-vere damage were observed inside the otherwise smooth weartrack.

The structural and the hardness analysis shows that the evolu-tion of the wear rate with increasing temperature can not be con-nected exclusively with either the oxidation or the considerablesoftening of the coatings. The increase of the wear rate in therange 300–400 �C for W70N30 is a typical example of such behav-iour, since both the hardness and the structure are almost iden-tical in this temperature range. The identification of thetungsten nitride wear mechanisms at elevated temperature is avery difficult task. Therefore, the discussion about the possiblewear mechanisms at elevated temperature is, in many points,speculative.

The observation of the wear track shows the moving of the po-sition of the adhered oxide layer with increasing temperature. Thethick layer localised in the wear track sides at room temperaturedisappears and a thin layer of tungsten oxide is formed in the cen-tre of the wear tracks at 300 �C. Moreover, the first cracks appear inall coatings at this temperature and they are located in the zonewith the oxide layer. Based on the results presented above, threemain wear mechanisms occur when the temperature is increased.The first dominant wear mechanism was the three-body wear witha thick tungsten trioxide tribolayer, typical for room temperatureand to some extent up to 300 �C. A further increase of the temper-ature (300–400 �C) leads to the vanishing of the adhered oxidelayer and the consequent crack formation. Finally, at 500 �C, thecracks propagate reaching the substrate and the rapid coatingdestruction is then observed, particularly in the centre of the weartrack. The evolution of the friction coefficient supports these pro-posed three stages. The initial value is the lowest, since the contactis represented mainly by the adhered tungsten oxide. As the ad-hered layer disappears with increasing temperature, the frictionrises, as it is observed. The friction stabilises in the range 300–400 �C, since there is no significant change observed in the weartracks of the individual coatings. The value of the friction coeffi-cient measured at 500 �C cannot be strictly considered as repre-senting the tungsten nitride (or oxide), since the coatings arepartially worn out.

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4. Concluding remarks

The tribological behaviour of a hard coating tested by pin-on-disc is often explained by the dominant influence of one of themechanical properties of the coatings. The improvement in thewear resistance is attributed to an increase of the hardness, thescratch-test resistance or the plasticity parameter H3/E2 amongmany others. Nevertheless, W–N coatings show a differentbehaviour. Although the hardness, adhesion/cohesion or mor-phology may influence the friction coefficient or the wear rateof the coating, none of these properties can be considered asdominant. Moreover, coatings exhibiting similar mechanicalproperties can show very different tribological behaviour. In thiscase, the type of the wear mechanism, together with the pres-ence of a wear debris interlayer forming a third-body, plays thedominant role in the tribological testing. In fact, the best wearresistance was achieved for the coatings with high N-contentexhibiting low hardness and low values of the scratch-test criticalloads.

With the exception of tribo-oxidiation, present in all testedcoated samples, the main wear mechanisms of low N-contentcoatings was abrasive wear, delamination and, in the particularcase of W88N12 sample, plastic deformation. On the contrary, thesliding of the coatings with high nitrogen content was predomi-nantly influenced by the formation of a third-body consisting ofvery small tungsten trioxide particles. Consequently, the mainwear mechanism was three-body abrasive wear with combinationwith polishing and mild wear. The improvement of the wear resis-tance in case of high N-content films was explained by the pres-ence of the protecting interlayer between the opposing surfacesin contact.

The selected W–N coating works well up to 200 �C due to for-mation of tungsten oxide tribolayer; further increase of the tem-perature prevents tribolayer formation leading to higher wearrates and coating damage.

Despite the existence of similar adhered tribolayer (coatingwith low W content) or similar dominant phase (high W content),the friction coefficient of oxides is higher than that of nitrides. Thisbehaviour has been attributed to the transfer of tungsten oxidematerial to the ball surface.

Acknowledgements

This work was supported by the Grant Agency of the Academyof Sciences of the Czech Republic through the projectKJB201240701 and by the Ministry of Education of the CzechRepublic (project MSM 6840770038).

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