Structure, mechanical and tribological characterisations of pulsed DC magnetron sputtered TiN-WSx composite coating

Vacuum 130 (2016) 93e104

Contents lists available at ScienceDirect

Vacuum journal homepage: www.elsevier.com/locate/vacuum

Structure, mechanical and tribological characterisations of pulsed DC magnetron sputtered TiN-WSx composite coating Tushar Banerjee*, A.K. Chattopadhyay Department of Mechanical Engineering, Indian Institute of Technology, Kharagpur, West Bengal 721302, India

a r t i c l e i n f o

a b s t r a c t

Article history: Received 15 March 2016 Received in revised form 30 April 2016 Accepted 2 May 2016 Available online 3 May 2016

Pulsed DC magnetron sputtered TiN-WSx composite coatings were investigated in terms of their structural, mechanical and tribological properties. The coatings were synthesized in Ar and N2 mixed environment from separate Ti and WS2 targets. Compositional variation was performed by systematically varying the Ti cathode current while keeping the same for WS2 cathode at a fixed value. EDS analysis of the films revealed that the cumulative (W þ S) content varied from 8.9 to 42.0 wt%. The films were, however, sulphur deficient, S/W at.% ratio being in the range of 0.6e0.78. GIXRD spectra indicated coexistence of (002) oriented WS2 phase along with polycrystalline TiN phase, as confirmed by HRTEM revealing presence of WS2 in nanocluster form within TiN columnar structure. FESEM micrographs showed decrease in agglomerated grain size and densification of columnar structure with increase in (W þ S) content. The adhesion also scaled with (W þ S) content up to 27.8 wt%. The nanohardness of the composite coatings, however, progressively decreased from 27.3 GPa to 17.3 GPa with increasing content of (W þ S). Significant reduction in friction coefficient and wear rate could also be recorded for the composite coatings, as compared to TiN, the best being obtained for the one containing 15.0 wt% of (W þ S). © 2016 Elsevier Ltd. All rights reserved.

Keywords: WS2 Structure Transmission electron microscopy XPS Adhesion Friction

1. Introduction MoS2 and WS2 are representatives of transition metal dichalcogenides which are explored in the form of solid lubricating coatings for vacuum and space applications [1]. The anisotropic crystal structure of MoS2 or WS2 consists of a sheet of metal (molybdenum or tungsten) atoms covalently bonded with sulphur layers on either side. The inter-lamellar sulphur layers are however weakly bonded through weaker Van der Waals forces which get easily sheared thereby providing low friction coefficient [2]. The lubricating property of pure MoS2 or WS2 coatings, in presence of moisture and humidity, however, deteriorates because of oxidation of the reactive edge sites as well as moisture penetration and entrapment through the coating thickness. Their inherent low hardness and relatively poor adhesion to the substrate also restricts their use in terrestrial applications. Addition of a range of metals like Ti, Cr, Zr, W, Ag, Ni, Ta etc to MoS2 or WS2 in tailored quantities, as investigated by several researchers, have led to some coinciding

* Corresponding author. E-mail address: [email protected] (T. Banerjee). http://dx.doi.org/10.1016/j.vacuum.2016.05.003 0042-207X/© 2016 Elsevier Ltd. All rights reserved.

observations like reduced crystallinity, denser microstructure, increased hardness and adhesion to the substrate, resulting in improved wear life [3e7]. Addition of non-metallic dopants, like C and N, especially to WS2, has also been explored and has been found to provide similar beneficial effects like metallic incorporation [8e10]. It has also been indicated that the non-metallic dopant N leaves the sliding interface in gaseous form, in contrast to its metallic counterparts, thereby allowing easier formation and functioning of tribo-film [10]. Industrial applications like cutting tool coatings for dry machining, however, calls for further improved mechanical properties of solid lubricant based coatings in order to prolong their useful service life. Though conventional hard coatings like TiN, TiAlN are frequently used, their crystal structure do not facilitate solid lubricating action and consequently they exhibit high friction coefficient in sliding contact against counter-bodies like tungsten carbide and bearing steel [4,11]. Sometimes the hard debris formed during the rubbing process further aggravates the magnitude of wear by participating in three-body abrasion with the mating parts. Therefore, there is an ongoing attempt to bring together the desirable properties like hardness and lubricity in a single coating by adopting the route of co-deposition of harder and lubricious

94

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

components. Some authors followed the route of co-sputtering of metals like Ti, Cr to form dispersions of ceramic phases like TiC, CrC along with W-S phase in a C matrix [12,13]. The hardness achieved, though higher as compared to undoped MoS2 or WS2, yet was limited to the maximum value of approximately 15 GPa [12]. Another route aimed towards retaining the ceramic phase as the dominant one and added tailored quantity of solid lubricant to them in order to improve their lubricious property. Representative members of such harder phases include TiN, TiSiN, CrN while MoS2 has been chosen as the softer counterpart in most of the studies [14e17]. Some investigations suggested that segregation of Mo and S occurred during deposition of TiN-MoS2 composite coating and no separate MoS2 phase could be detected through XRD analysis [14,18]. Mo and S, in elemental form, was speculated to be incorporated into the cubic TiN lattice or segregated at the grain boundaries. Retention of high hardness of the composite at par with that of TiN, also indicated towards absence of any softer phase in the as deposited coating [14]. However, the eventual reduction in friction coefficient and wear rate of the composite, as compared to pure TiN was attributed to friction induced recombination of Mo and S during the sliding process [14,18]. Some reports on pulsed magnetron sputtered TiN-MoS2 composite coating, however, revealed that MoS2 can exist as a separate phase along with TiN in the composite and such coatings can possibly perform better in tribological tests because of ready availability of lubricating phase in the coating [15,19]. Despite exhibiting similar crystal structures, the research on WS2 has been limited compared to that of MoS2. WS2 excels MoS2 in terms of mass loss at increasing temperatures and provides effective lubrication at temperatures approximately 200 C higher as compared to that of MoS2 [1,20]. A composite coating of CrNWS2, synthesized by unbalanced magnetron sputtering, showed denser coating morphology and accompanied improvement in hardness, adhesion as well as elastic recovery ability as compared to pure WS2 [21]. Though the friction coefficient was moderately higher (m ~ 0.2e0.24), it significantly outperformed pure WS2 in terms of tribological endurance. However, comparing such ceramic based composites with that of the pure ceramic counterparts in terms of wear resistance property may be of relatively greater interest from the industrial viewpoint of dry machining. TiN is a well known hard coating which has been satisfactorily serving the cutting tool industry for decades. The present work therefore aims at studying the composition e structure e property e tribological performance correlation of hard-lubricious TiN-WSx composite coating. Comparative analysis of different compositions of the composite has been done with respect to a TiN coating. The nature of existence of WS2 inside the TiN matrix, which is not thoroughly investigated in the literature, has also been given attention in this work. The coatings were deposited by pulsed DC closed field unbalanced magnetron sputtering technique along with pulsed substrate biasing. Such hard-lubricious coatings have potential for application as cutting tool coatings for dry machining as well as other mechanical components like bearings, gears etc. 2. Experimental methods and conditions The coatings were deposited by a dual cathode closed field unbalanced magnetron sputtering system [4] on IS C40 (AISI 1040) steel discs (25 mm diameter and 8 mm thickness), M2 grade HSS blocks (10 mm  10 mm  20 mm), stainless steel foils and ISO K10 grade cemented carbide turning inserts. The coated discs were used for tribological tests and the coated HSS blocks, stainless steel foils and cemented carbide inserts were used for other physical and mechanical characterisations. Before deposition, the discs and the

HSS blocks were polished to obtain a surface finish of 50 nm or better. The deposition system contained one Ti target and one WS2 target (with the dimensions of 254 mm  127 mm  12 mm and 99.9% purity) placed vertically opposite to each other. The substrates were mounted on a substrate holder assembly placed in between the two targets and two-fold rotation was imparted to them to ensure uniform coating coverage on all regions. The targets as well as the substrates were connected to individual pulsed DC power supply units (Advanced Energy Pinnacle Plus) having maximum power capacity of 5 kW each. The targets were operated in ‘current’ regulation mode while the substrates were operated in ‘voltage’ regulation mode. The chamber was initially pumped down to a base vacuum of approximately 1  103 Pa and heated to 200 C, after which Ar was flown in at 15 sccm through a mass flow controller (MKS instruments). In order to reduce water vapour content inside the chamber, Ti sputtering was carried out, before actual coating deposition, with the target shutter in closed condition. This was followed by Ti-ion etching of the substrates at a pulsed (35 kHz, 90% ON) substrate bias of 400 V. After that a Ti layer of approximately 100 nm was deposited to promote filmsubstrate adhesion. In order to bring the compositional variation for different specimens, the Ti cathode current and corresponding N2 flow rate was varied while keeping the WS2 cathode current at a fixed value, since increasing the cathode current of WS2/MoS2 leads to severe loss of sulphur due to increasingly energetic ion bombardment on the cathode surface [11,16]. The parametric variations for the different specimens are provided in Table 1. For all the specimens, a pulsed (35 kHz, 90% ON) substrate bias of 50 V was maintained during the entire deposition duration. All depositions were carried out under a process pressure of 0.3 Pa. Surface morphology and cross sectional features of the coatings were imaged with a field emission scanning electron microscope (Carl Zeiss, Supra 40) and the chemical compositions were studied with an energy dispersive X-ray microanalysis system attached with the FESEM. The crystalline structure of the coatings were studied with X-ray diffraction in grazing incidence mode (Philips, PANalytical PW 3040/60 X'Pert PRO), with Cu Ka radiation (45 kV and 40 mA) of 0.15418 nm wavelength at an incident angle of 2 . High resolution transmission electron microscopy (JEOL JEM 2100 instrument with a resolution of 0.14 nm) was employed to study the cross sectional structure of the composite coating. The cross section of the sample (deposited on stainless steel foil) for TEM was prepared through focussed ion-beam technique. X-ray photoelectron spectroscopy (XPS) analysis (Model PHI 5000 Versa Probe II, ULVAC-PHI, Inc.) was carried out with Al Ka radiation (1486.6 eV) at 15 kV. The XPS data was calibrated with reference to C 1s peak at 284.5 eV. It may be mentioned here that the sample was transported in a vacuum desiccator to the XPS analysis chamber just after deposition to minimize surface oxidation. No surface etching was done before XPS analysis of the sample, since WS2 based coatings are susceptible to sputter damage. The adhesion of the coatings, deposited on polished HSS substrates were evaluated with a scratch adhesion tester (TR-101 M5, DUCOM) equipped with a Rockwell C diamond stylus with 200 mm tip radius. The stylus was drawn over the coated surfaces with an increasing normal load in the range of 10e80 N at loading rate of 5 N/mm and velocity of 0.2 mm/s. The point of sharp increase of the friction force, as obtained through the associated software WINDUCOM 2004, indicated detachment of the coating and substrate exposure from the centre of the scratch track and the corresponding normal load was recorded as the critical load (Lc3). However, for confirming the critical load values and studying the nature of failure, the scratch tracks were later visualized through an inverted metallurgical microscope (ZEISS AxioCam ERc5s). The adhesion of the different coatings deposited on ISO K10 grade cemented carbide turning

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

95

Table 1 Deposition details of different coatings. Ti cathode current

S0 S1 S2 S3 S4

6.5 6.5 5.0 3.5 2.0

A A A A A

(35 (35 (35 (35 (35

kHz, kHz, kHz, kHz, kHz,

60% 60% 60% 60% 60%

ON) ON) ON) ON) ON)

inserts were further assessed by indentation adhesion test with a Rockwell C diamond indenter under a normal load of 600 N. The hardness of the coatings deposited on polished HSS substrates were measured with a nanoindenter (TI950 Tribonanoindenter, Hysitron Inc., Minnesota, MN, USA) using a maximum load of 8 mN with loading and unloading time of 10 s each. The hardness values were then calculated following Oliver Pharr method [22]. The tribological performance of the coatings was evaluated with a tribometer (CSM instruments) in pin-on-disc configuration. A WC ball of 6 mm diameter was used as the counter body with normal load of 5 N and sliding velocity of 10 cm/s. All the tests were carried out in room temperature of 25e30 C and relative humidity of 50 ± 5%. Raman spectroscopy (Model: T64000, Make: Jobin Yvon Horiba, France) with an Arþ laser having excitation wavelength of 488 nm at 80 mW power was employed for identifying the tribo-chemical phases after tribological tests. The cross-sectional profiles of the wear tracks were later recorded with a surface profilometer (Taylor- Hobson, Surtronic 3þ) for estimating the wear coefficients. 3. Results and discussion 3.1. Chemical composition The chemical composition of the coatings deposited on HSS substrates, as obtained from EDS analysis, are presented in Table 2. It can be observed that, decreasing the Ti cathode current from 6.5 to 2.0 A resulted in increase in (W þ S) content in the coatings from 8.9 to 42.0 wt%, expectedly because of reduction in deposition rate of TiN. The coatings, however, were sulphur deficient exhibiting overall S/W at.% ratio in the range of 0.6e0.78 as compared to the stoichiometric value of 2 for WS2. Similar values of S/Mo at.% ratio were also reported for magnetron sputtered composite coatings like TiN-MoSx, TiSiN-MoSx, TiB2-MoS2 [14,16,23]. Sulphur deficiency in sputter deposited MoS2 and WS2 based films is a widely reported phenomenon and has been attributed to a number of factors like different sputtering rates of molybdenum/tungsten as compared to sulphur from MoS2/WS2 target, different degree of scattering in the gaseous phase, presence of reactive gases like N2 and CH4 in the sputtering chamber, their different sticking coefficients during film growth and finally preferential resputtering of lighter sulphur atoms as compared to molybdenum/tungsten atoms from the growing film by bombardment of energetic particles [24,25]. The application of pulsed DC to the targets as well as to the substrates along with closed field unbalanced configuration of the magnetrons in the present work also augmented the degree of

Table 2 EDS analysis of the coatings.

N2 flow rate

WS2 cathode current

13.5 sccm 13.5 sccm 11.5 sccm 10.0 sccm 8.5 sccm

e 0.2 0.2 0.2 0.2

(W þ S) wt%

S/W at.% ratio

S0 S1 S2 S3 S4

e 8.9 15.0 27.8 42.0

e 0.60 0.71 0.76 0.78

(100 (100 (100 (100

kHz, kHz, kHz, kHz,

60% 60% 60% 60%

ON) ON) ON) ON)

3.2. Structure, morphology and chemical bonding 3.2.1. Structure Fig. 1 shows the GIXRD spectra of the different coatings deposited on M2 grade HSS substrates. The TiN coating (specimen S0) showed a polycrystalline structure with no particular preferred orientation. During initial stages of TiN growth, the (111) orientation becomes kinetically most favoured because of lattice similarity between the mono-atomic (Ti or N) layers and the Ti adhesion layer [28,29]. However, during coating growth, continuous ion bombardment causes an increase in the substrate temperature which promotes atomic mobility in the deposits. This results in evolution of other crystallographic planes like (200), (220), (311) because of mass transfer from the initially predominant (111) oriented regions [28]. The application of substrate bias in pulsed mode along with closed field unbalanced configuration of the magnetrons in the present work, also enhanced such ion bombardment leading to a mixed orientation of the TiN film (S0) [29]. However, on addition of WSx to TiN (for S1), the orientation changed to strong (111). Such behaviour has also been reported for magnetron sputtered TiN-MoS2 and TiSiN-MoS2 composite coatings [14,16]. This phenomenon may be attributed to the restriction imposed to the atomic mobility of Ti and N from (111) regions to the other regions because of presence of WSx in the deposits. The GIXRD spectra corresponding to specimens S2eS4, exhibited a progressive weakening of the (111) orientation, along with gradual evolution of (200), (220), (311), (222). It may be noted here that, because of

a- WS2 (002) d- TiN (220) b- TiN (111) e- TiN (311) c- TiN (200) f- TiN (222)

1400 1200

S4

1000 S3

800 600

S2

400

S1

a

200

Specimen code

A A A A

ion-bombardment on the growing film thereby promoting such sulphur loss [4,24,26]. However, it has been demonstrated that despite providing lower S/Mo at.% ratio, proper selection of pulsing frequency and duty cycle during sputtering, ensures a basally oriented film, which is favourable for tribological performance [25e27].

Intensity (a.u.)

Specimen code

d

c

b

S0

e

f

0 0

10

20

30

40

50

60

70

2θ Fig. 1. GIXRD spectra of the different coatings.

80

90

96

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

progressive reduction in Ti cathode currents for the specimens S2eS4 (as compared to S1), the corresponding deposition durations were extended in order to achieve coating thicknesses of the same order for all the specimens (as discussed in section 3.2.2. Study of morphology and cross section of the coatings). Such extended deposition durations resulted in prolonged ion-bombardment during film growth thereby allowing more time for atomic rearrangement of Ti and N through surface mobility. This in turn caused a weakening of (111) orientation and evolution of other crystallographic planes. The spectra corresponding to the composite coatings S1eS4 showed coexistence of (002) basal plane of WS2 at approximately 2q ¼ 14 in the form of a broad hump along with crystalline planes of TiN which indicates the presence of WS2 in nanocrystalline form in the composite coating [15,19]. Some literatures, however, suggested segregation of Mo and S in the grain boundaries of TiN and XRD analysis could not detect the presence of any separate MoS2 phase [14,18]. With increase in (W þ S) content, up to coating S3, the peaks of TiN gradually shifted to lower angles, indicating increase in the lattice parameter of TiN owing to incorporation of tensile stresses [15,19]. In order to elucidate the nanostructure of TiN-WSx composite coating, high resolution cross sectional transmission electron microscopy was carried out. Fig. 2(a) shows the cross sectional HRTEM image of coating S2 deposited on stainless steel foil. The image features a thin Ti adhesion layer of approximately 100 nm thickness followed by a graded interlayer leading to TiN which grew with a columnar structure. Within the TiN columnar structure, some isolated dark clusters could be identified, one of which is shown through progressively higher magnifications in Fig. 2(b)e(c). Fig. 2(d) shows a detail view of the circular region marked in

Fig. 2(c), resolving the directional planes of the crystallite. The corresponding d-spacing was measured to be approximately 0.6 nm (6 Å) which is in good agreement with that of (002) planes of WS2 [30]. Moreover, it is worth noting that the crystallite is oriented with its basal plane perpendicular to the coating growth direction (through comparison of the images in Fig. 2(a)e(d)) which is favourable for providing low friction by inducing easy slip. However, such observation of a localized WS2 crystallite and its orientation within the TiN matrix is stochastic in nature and may be argued on the issue of uniformity of its distribution. Nevertheless, the HRTEM images corroborate the results obtained from GIXRD analysis (Fig. 1) regarding presence of separate WS2 phase and confirm its presence in nanocluster form, at least in some quantity, within the TiN matrix under the current deposition conditions. 3.2.2. Study of morphology and cross section of the coatings The FESEM micrographs of surface morphologies and cross sections of the coatings deposited on M2 grade HSS substrates are presented in Fig. 3. The morphology of pure TiN coating showed an elongated granular appearance [31]. With addition of WSx (specimen S1), the morphology exhibited agglomerated granular structure. Further increase in (W þ S) content, through reduction of Ti cathode current, resulted in progressive refinement of the agglomerated grain size as observed for specimens S2eS4. This phenomenon may be attributed to the competitive growth of TiN and WSx in the composite coating, which hindered the growth of individual grains [32]. Similar observation regarding grain refinement with increase in content of softer phase was also made for TiN-MoS2, CrN-WS2 composite coatings [15,21]. The FESEM micrographs of the corresponding fractured cross

Fig. 2. Cross sectional HRTEM of coating S2.

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

Fig. 3. Surface morphology and cross section of the different coatings.

97

98

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

sections of the coatings are also presented in Fig. 3. The thicknesses of all the coatings were of the order of 2 mm. The fractographs feature a thin Ti adhesion layer of approximately 100 nm deposited initially on the substrates in order to promote film substrate adhesion. A columnar structure was observed for all the coatings which is typical of sputtered transition metal nitrides [19,21]. However, the overall width of the columns reduced with progressive increase in (W þ S) content and the structure became more compact which correlates with the observations made for change in surface morphology of the coatings.

3.2.3. Chemical bonding To identify the chemical bonding states of Ti, W, N and S, XPS analysis was carried out for composite coating S2 (containing 15.0 wt% of (W þ S)) and the resultant spectra are presented in Fig. 4. The Ti 2p spectra (Fig. 4(a)) could be deconvoluted to two doublets. The peaks located at 455.7 and 461.7 eV correspond to Ti 2p3/2 and Ti 2p1/2 respectively for TiN while those at 457.9 and 464.1 eV correspond to that of TiO2 [33]. To obtain a good fit for the W 4f spectra, contribution from three doublets was necessary as shown in Fig. 4(b). The peaks at 32.0 and 34.2 eV correspond to W 4f7/2 and W 4f5/2 respectively for W-S [21,34], at 33.7 and 35.4 eV for W-N [33,34] while at 37.0 and 38.2 eV for W-O [34]. The higher electro negativity of N as compared to S (3.04 and 2.58 respectively), resulted in appearance of W-N peaks at higher binding energy as compared to that of W-S. The deconvoluted N 1s spectra (Fig. 4(c)) consists of peaks located at 396.8 eV which corresponds to TiN, W-N and at 398 eV for N-oxide [33]. The S 2p spectra (Fig. 4(d)) could be deconvoluted to a doublet consisting of S 2p3/2 and S 2p1/2 peaks located at binding energy values of 161.3 and

(a)

3.3. Scratch adhesion test Typical variation of friction forces with increase in normal load during scratch adhesion test of different coatings are presented in Fig. 5(a). The corresponding critical load (Lc3) values are presented in Fig. 5(b). The critical load (Lc3) values were considered as the normal load at which a sudden increase in the friction force occurred, as confirmed through removal of the coating and exposure of the substrate from the middle of the scratch track. Typical appearances of scratch tracks at the failure zone are represented in Fig. 6. The TiN coating (S0) showed an average critical load of 50 N. The corresponding failure (Fig. 6(a)) zone showed flaking on the two sides of the scratch tracks which is indicative of the brittleness of the TiN film. The stored energy at the tip of the moving stylus was utilized for opening up cracks on the two sides of the scratch track [36]. With increase in (W þ S) content up to 27.8 wt% (corresponding to specimen S3), the critical load increased up to 68 N after which a steep drop was observed for coating S4. The improvement may be attributed to denser coating structure owing to reduced agglomerated grain size as observed through the FESEM micrographs (Fig. 3). The brittle fractures at the sides of the scratch tracks, as observed for TiN (S0), significantly reduced for specimen S1 containing 8.9 wt% of (W þ S) and almost got eliminated for coatings S2 and S3, containing 15.0 and 27.8 wt% of (W þ S) respectively. This indicates the capability of the composite coatings

(b)

Ti 2p3/2 455.7 eV

W 4f5/2 34.2 eV

Ti 2p1/2 461.7 eV

Intensity (a.u.)

Ti 2p3/2 457.9 eV

454

456

458

460

462

464

466

W 4f7/2 33.7 eV

W 4f7/2 32.0 eV

Intensity (a.u.)

Ti 2p1/2 464.1 eV

452

162.8 eV respectively which corresponds to S in WS2 [12,21]. The position of S 2p3/2 and S 2p1/2 peaks at lower binding energy values as compared to literature [35] indicates towards sulphur deficiency in the films.

W 4f5/2 35.4 eV W 4f7/2 37.0 eV

468

28

30

32

Binding energy (eV) (c)

34

36

38

40

W 4f5/2 38.2 eV

42

44

Binding energy (eV) S 2p3/2 161.3 eV

(d)

396.8 eV

Intensity (a.u.)

Intensity (a.u.)

S 2p1/2 162.8 eV

398.0 eV

392

394

396

398

Binding energy (eV)

400

402

158

159

160

161

162

163

Binding energy (eV)

Fig. 4. XPS spectra of (a) Ti 2p, (b) W 4f, (c) N 1s and (d) S 2p for coating S2.

164

165

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

(a)

to accommodate the stresses induced by the moving stylus and may be attributed to the presence of softer phase in the coatings which hindered the propagation of cracks. Further increase in (W þ S) content to 42.0 wt%, however, reduced the coating hardness (as discussed in section 3.5 Nanoindentation test), thereby allowing easier penetration of the scratch indenter through the coating, and detaching it off the substrate at lesser value of normal load.

80

Force (N)

60 Normal Load

Friction Forces

40

3.4. Indentation adhesion test

S0 S1 S2 S3 S4

20

70 (b)

The adhesion property of the different coatings deposited on ISO K10 grade WC turning inserts was further studied through indentation adhesion test. Fig. 7 shows the optical images of the indents done on the coatings under normal load of 600 N. Coating delamination along the edges of the indent is evident for TiN (S0) which is negligible for all the composite coatings. The presence of softer phases in the coatings successfully restricted the propagation of radial cracks which led to such improved property. Such observations coincided with that for scratch adhesion test (Fig. 6) and imply that both toughness and adhesion of the coatings improved with addition of WSx to TiN.

60

3.5. Nanoindentation test

50

Representative loading-unloading plots showing the average indentation depths for each specimen are shown in Fig. 8(a). For all the specimens, the indentation depth was limited within 1/10th of the coating thickness in order to eliminate the effect of substrate hardness. The variation of nanohardness with increase in (W þ S) content is shown in Fig. 8(b). The TiN coating (specimen S0) exhibited an average nanohardness of approximately 26 GPa which lies within the range of mostly reported values for magnetron sputtered TiN [14,31,32]. Addition of WSx to TiN initially led to an increase in hardness (for specimen S1) which can be explained by the compaction of the coating structure because of competitive growth of TiN and WSx. The hardness, however, progressively decreased with further addition of (W þ S) because of formation of extended softer regions in the composite. It may be noted here, that the decrease in nanohardness coincided with the decrease in

0 0

2

4

6

8

10

12

14

16

Scratch Length (mm)

Critical Load (N)

99

40 30 20 10 0 S0

S1

S2

S3

S4

Fig. 5. (a) Representative plots for variation of friction forces with increase in normal load during scratch adhesion test (b) Critical load (Lc3) for different coatings.

Fig. 6. Optical images of the failure zones of different coatings in scratch adhesion test: (a) S0, (b) S1, (c) S2, (d) S3, (e) S4.

100

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

Fig. 7. Optical images of indents on different coatings in indentation adhesion test: (a) S0, (b) S1, (c) S2, (d) S3, (e) S4.

8

(a)

S1 S0 S2 S3

S4

7

FESEM micrographs (Fig. 3). Similar observation in hardness reduction with reduction in agglomerated grain size was also reported for magnetron sputtered TiN-MoSx, CrN-WS2, CrB2-MoS2 composite coatings [15,21,37].

Load (mN)

6 5

3.6. Friction coefficient and wear behaviour

4

The variation of friction coefficients with sliding distance for the different coatings deposited on polished C40 steel discs is presented in Fig. 9. The TiN coating (specimen S0) showed a steady state friction coefficient of approximately 0.7 which conforms to the mostly reported values for TiN [14,38,39]. High fluctuations were registered for the entire duration of the test. In comparison, lesser values of friction coefficients were registered for all the composite coatings (specimens S1eS4). Specimens S2 and S3 (containing approximately 15.0 and 27.8 wt% of (W þ S) respectively) performed superior in the tribological test, exhibiting steady state values in the range of 0.25e0.35. Significant reduction in fluctuations of friction coefficient could also be observed, especially

3 2 1 0 0

20

40

60

80

100

120

Indentation Depth (nm)

30

(b)

1.0 0.9 0.8

20

Coefficient of friction

Nanohardness (GPa)

25

15 10 5

0.7

S0

0.6

S4

0.5

S1

0.4

S3 S2

0.3 0.2

0 S0

S1

S2

S3

S4

Fig. 8. (a) Representative loading-unloading curves obtained from nanoindentation measurements for different coatings (b) Nanohardness of different coatings.

0.1 0.0 0

100

200

300

400

500

600

700

Sliding Distance (m) agglomerated grain size of the coatings as observed from the

Fig. 9. Variation of coefficient of friction with sliding distance for different coatings.

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

for specimen S2, as compared to that of TiN coating (S0). It may be recalled here that in the range of (W þ S) content of 15.0e27.8 wt%, the critical loads of adhesion (Lc3) and nanohardness (corresponding to specimens S2 and S3) were of the same order (Figs. 5(b) and Fig. 8(b)), which possibly led to such comparable values of friction coefficients. Further increase in (W þ S) content to approximately 42.0 wt%, however, led to an increase in friction coefficient along with fluctuating nature. The deterioration in mechanical properties like adhesion and nanohardness, were responsible for such poorer tribological performance at higher end of (W þ S) content. Nevertheless, the still lower average value of friction coefficient of specimen S4, as compared to that of TiN coating (S0), may be attributed to the presence of WSx in the wear track which could sustain the solid lubricating action to some extent. Typical appearances of wear tracks of the different coatings after 700 m of dry sliding under a normal load of 5 N are shown in Fig. 10. The widths of the wear tracks decreased with increase in (W þ S) content up to coating S2 (containing 15.0 wt% of (W þ S)), which also exhibited the lowest steady state friction coefficient of

101

approximately 0.25. Further addition, however deteriorated the wear resistance as evident from the increased wear track width, particularly for coating S5 with highest content of (W þ S). The wear track corresponding to TiN coating (S0) featured several fractured and eroded regions amidst some parallel running furrows (Fig. 10(a)). Prolonged rubbing against hard WC ball counter body produced harder debris from both sides which participated in three-body abrasion with the tribo-pair resulting in coating removal in the form of chunks. A higher magnification image in the vicinity of such an eroded area clearly reveals the cracks in the coating (Fig. 11(a)). EDS spectra acquired from the region marked A (Fig. 11(d)) showed prominent presence of Fe as compared to that from region B (Fig. 11(e)), confirming substrate exposure due to such chunk removal. Peaks of W in the EDS spectra (Fig. 11(d) and (e)) originated from wear products of WC ball counter body. In comparison, the wear track morphology of coating S2 appeared much smoother (Fig. 10(c)). Higher magnification image of the same (Fig. 11(b)) showed some smeared and plastically deformed regions which are in contrast to the fractured and delaminated regions observed for TiN coating (S0) (Fig. 11(a)). This

Fig. 10. Representative SEM images of the wear tracks of coating (a) S0 (TiN), (b) S1 (8.9 wt% (W þ S)), (c) S2 (15.0 wt% (W þ S)), (d) S3 (27.8 wt% (W þ S)) and (e) S4 (42.0 wt% (W þ S)).

102

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

Fig. 11. High magnification SEM images of wear tracks of coating (a) S0 (TiN), (b) S2 (15.0 wt% (W þ S)) (c) S4 (42.0 wt% (W þ S)) and (d)e(g) EDS analysis of the wear tracks.

observation could be correlated with the results obtaining through scratch adhesion test (Fig. 6(a) and (c)) where minimal lateral flaking was observed for coating S2 as compared to TiN (S0). EDS spectra (Fig. 11(f)) acquired from the region marked C showed presence of coating elements like Ti, N, W and S indicating that the coating was still retained after 700 m of sliding. The wear track of coating S4, in addition to increased width, also featured some delaminated regions as observed from the higher magnification image (Fig. 11(c)). Presence of Fe in the EDS spectra (Fig. 11(g)) acquired from the region marked D indicates thinning of the coating and partial exposure of the substrate in those regions. In order to further investigate formation of tribo-chemical

phases, Raman spectra were acquired from different spots near the central region of wear track of coating S2 and presented in Fig. 12. The absence of clearly discernible peaks for WS2 at approximate positions of 350 cm1 and 415 cm1 [4,30], indicate that the structural order of the same was low, or the tribo-layer, if formed, was too thin to be detected [40]. Similar Raman spectra, with no noticeable presence of WS2 phase were also reported for wear tracks of composite coatings like W-S-C/Cr, W-S-N, exhibiting friction coefficient in the range of 0.1e0.3 while sliding in humid air [13,40]. The super-lubricious property (m << 0.1) demonstrated by WS2 based coatings, has however, been attributed to sliding induced formation of a well-ordered (002) oriented tribo-layer on

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

103

4. Conclusions

Intensity (a.u.)

WS2 WS2

300

400

500

600

700 -1 Raman Shift (cm )

800

900

Fig. 12. Raman analysis of wear track of coating S2.

the wear track and counter-body [10,30]. The ceramic based composite coatings investigated in the present work, in contrast, have major content of TiN with dispersions of WSx grains; therefore, such well-aligned tribolayer formation during sliding is not expected in this case. Rather, continuous wear-induced exposure of soft lubricious WSx is responsible for reduced friction coefficient (m ~ 0.2e0.3) as compared to that of TiN (m ~ 0.7), but higher than that for super-lubricious WS2 based coatings (m << 0.1). Nevertheless, it is worth noting here that, such ceramic based composites have a major advantage of hardness of the order of 20e25 GPa, which is notably higher than that of metal or non-metal doped WS2 based coatings (5e10 GPa) [4e10]. Such higher hardness is important for withstanding plastic deformation and abrasive wear in particular, typically encountered during mechanical applications like dry machining. The wear coefficients of the different coatings after 700 m of sliding are presented in Fig. 13. With progressive addition of (W þ S), the wear coefficient decreased up to specimen S2, which exhibited the lowest steady state friction coefficient with less fluctuations. Further increase in (W þ S) content however, deteriorated the wear resistance, especially for coating S4, because of fall in adhesion and nanohardness.

Wear coefficient (x 10

-16

3 ) (m /N-m)

100

80

60

40

20

0 S0

S1

S2

S3

Fig. 13. Wear coefficient of different coatings.

S4

It was possible to synthesize hard-lubricious TiN-WSx composite coatings by pulsed DC closed field unbalanced magnetron sputtering. The (W þ S) content in the coatings were varied from approximately 8.9e42.0 wt% for different specimens, by progressively reducing the Ti cathode current and N2 flow rate, while keeping the WS2 cathode current to a fixed value. The GIXRD spectra of the coatings indicated presence of (002) oriented WS2 phase along with polycrystalline TiN, which was confirmed by cross-sectional HRTEM revealing WS2 nanoclusters embedded within TiN columnar structure. The agglomerated grain size, observed through FESEM, progressively became finer with increasing content of (W þ S). The critical load of adhesion also improved with (W þ S) content up to approximately 27.8 wt%; further increment, however, led to deterioration of the same. The brittle flakings on the sides of the scratch tracks, as observed for TiN were almost eliminated for the composite coatings. The nanohardness of the coatings progressively reduced with increase in (W þ S) content. The composite coating containing approximately 15.0 wt% on (W þ S) performed the best in the tribological test, exhibiting much lower friction coefficient and wear rate as compared to TiN coating. Acknowledgements The authors gratefully acknowledge the XPS facility (Grant No. SR/FST/PSII-007/2010, Dt. 29.10.2010) provided by Department of Physics, Indian Institute of Technology Kharagpur. The assistance provided by Mr. Uttam Kumar Ghara in operating the XPS system and Mr. Velalam Abhilash and Mr. Shashank Pashikanti during coating deposition is also acknowledged. References [1] W.A. Brainard, The Thermal Stability and Friction of the Disulfides, Diselenides and Ditellurides of Molybdenum and Tungsten in Vacuum (109 to 106 Torr), NASA TN D-5141, 1969. [2] W.O. Winer, Molybdenum disulfide as a lubricant: a review of the fundamental knowledge, Wear 10 (1967) 422e452. [3] N.M. Renevier, V.C. Fox, D.G. Teer, J. Hampshire, Coating characteristics and tribological properties of sputter-deposited MoS2 metal composite coatings deposited by closed field unbalanced magnetron sputter ion plating, Surf. Coat. Technol. 127 (2000) 24e37. [4] T. Banerjee, A.K. Chattopadhyay, Structural, mechanical and tribological properties of WS2-Ti composite coating with and without hard under layer of TiN, Surf. Coat. Technol. 258 (2014) 849e860. [5] Q. Wang, J.P. Tu, S.C. Zhang, D.M. Lai, S.M. Peng, B. Gu, Effect of Ag content on microstructure and tribological performance of WS2eAg composite films, Surf. Coat. Technol. 201 (2006) 1666e1670. [6] S. Xu, X. Gao, M. Hu, J. Sun, D. Jiang, F. Zhou, W. Liu, L. Weng, Nanostructured WS2eNi composite films for improved oxidation, resistance and tribological performance, Appl. Surf. Sci. 288 (2014) 15e25. [7] X. Liu, G.J. Ma, G. Sun, Y.P. Duan, S.H. Liu, MoSx-Ta composite coatings on steel by d.c magnetron sputtering, Vacuum 89 (2013) 203e208. [8] A. Nossa, A. Cavaleiro, N.J.M. Carvalho, B.J. Kooi, J.Th.M. De Hosson, On the microstructure of tungsten disulfide films alloyed with carbon and nitrogen, Thin Solid Films 484 (2005) 389e395. [9] T. Polcar, M. Evaristo, A. Cavaleiro, The tribological behavior of WeSeC films in pin-on-disk testing at elevated temperature, Vacuum 81 (2007) 1439e1442. [10] F. Gustavsson, S. Jacobson, A. Cavaleiro, T. Polcar, Ultra-low friction WeSeN solid lubricant coating, Surf. Coat. Technol. 232 (2013) 541e548. [11] T. Banerjee, A.K. Chattopadhyay, Structural, mechanical and tribological properties of pulsed DC magnetron sputtered TiNeWSx/TiN bilayer coating, Surf. Coat. Technol. 282 (2015) 24e35. €rhammar, F. Gustavsson, T. Kubart, T. Nyberg, [12] J. Sundberg, H. Nyberg, E. Sa S. Jacobson, U. Jansson, Influence of Ti addition on the structure and properties of low-friction WeSeC coatings, Surf. Coat. Technol. 232 (2013) 340e348. [13] T. Polcar, F. Gustavsson, T. Thersleff, S. Jacobson, A. Cavaleiro, Complex frictional analysis of self-lubricant W-S-C/Cr coating, Faraday Discuss 156 (2012) 383. [14] R. Goller, P. Torri, M.A. Baker, R. Gilmore, W. Gissler, The deposition of lowfriction TiNeMoSx hard coatings by a combined arc evaporation and

104

T. Banerjee, A.K. Chattopadhyay / Vacuum 130 (2016) 93e104

magnetron sputter process, Surf. Coat. Technol. 120e121 (1999) 453e457. [15] S. Gangopadhyay, R. Acharya, A.K. Chattopadhyay, S. Paul, Composition and structureeproperty relationship of low friction, wear resistant TiNeMoSx composite coating deposited by pulsed closed-field unbalanced magnetron sputtering, Surf. Coat. Technol. 203 (2009) 1565e1572. [16] X.Z. Ding, X.T. Zeng, T. Goto, Unbalanced magnetron sputtered TieSieN: MoSx composite coatings for improvement of tribological properties, Surf. Coat. Technol. 198 (2005) 432e436. [17] S.K. Kim, B.C. Cha, Deposition of CrNeMoS2 thin films by D.C. magnetron sputtering, Surf. Coat. Technol. 188e189 (2004) 174e178. [18] M. Rahman, J. Haider, D.P. Dowling, P. Duggan, M.S.J. Hashmi, Deposition of magnetron sputtered TiNþMoSx coating with TieTiN graded interlayer, Surf. Coat. Technol. 200 (2005) 1071e1075. [19] G. Strapasson, P.C. Badin, G.V. Soares, G. Machado, C.A. Figueroa, R. Hubler, A.L. Gasparin, I.J.R. Baumvol, C. Aguzzoli, E.K. Tentardini, Structure, composition and mechanical characterization of dc sputtered TiN-MoS2 nanocomposite thin films, Surf. Coat. Technol. 205 (2011) 3810e3815. [20] H.E. Sliney, Solid lubricant coatings for high temperatures e a review, Tribol. Int. 15 (1982) 303e315. [21] B. Deepthi, H.C. Barshilia, K.S. Rajam, M.S. Konchady, D.M. Pai, J. Sankar, Structural, mechanical and tribological investigations of sputter deposited CrNeWS2 nanocomposite solid lubricant coatings, Tribol. Int. 44 (2011) 1844e1851. [22] W.C. Oliver, G.M. Pharr, An improved technique for determining hardness and elastic modulus using load and displacement sensing indentation experiments, J. Mater. Res. 7 (1992) 1564e1583. [23] R. Gilmore, M.A. Baker, P.N. Gibson, W. Gissler, Preparation and characterisation of low-friction TiB2-based coatings by incorporation of C or MoS2, Surf. Coat. Technol. 105 (1998) 45e50. €rhammar, E. Strandberg, J. Sundberg, H. Nyberg, T. Kubart, S. Jacobson, [24] E. Sa U. Jansson, T. Nyberg, Mechanisms for compositional variations of coatings sputtered from a WS2 target, Surf. Coat. Technol. 252 (2014) 186e190. [25] J. Wang, W. Lauwerens, E. Wieers, L.M. Stals, J. He, J.P. Celis, Structure and tribological properties of MoSx coatings prepared by bipolar DC magnetron sputtering, Surf. Coat. Technol. 139 (2001) 143e152. [26] W. Lauwerens, J. Wang, J. Navratil, E. Wieers, J. D'haen, L.M. Stals, J.P. Celis, Y. Bruynseraede, Humidity resistant MoSx films prepared by pulsed magnetron sputtering, Surf. Coat. Technol. 131 (2000) 216e221. [27] C. Muratore, A.A. Voevodin, Control of molybdenum disulfide basal plane orientation during coating growth in pulsed magnetron sputtering discharges,

Thin Solid Films 517 (2009) 5605e5610. [28] I. Iordanova, P.J. Kelly, R. Mirchev, V. Antonov, Crystallography of magnetron sputtered TiN coatings on steel substrates, Vacuum 81 (2007) 830e842. [29] I. Iordanova, P.J. Kelly, M. Burova, A. Andreeva, B. Stefanova, Influence of thickness on the crystallography and surface topography of TiN nano-films deposited by reactive DC and pulsed magnetron sputtering, Thin Solid Films 520 (2012) 5333e5339. [30] T.W. Scharf, A. Rajendran, R. Banerjee, F. Sequeda, Growth, structure and friction behavior of titanium doped tungsten disulphide (Ti-WS2) nanocomposite thin films, Thin Solid Films 517 (2009) 5666e5675. [31] D. Bhaduri, A. Ghosh, S. Gangopadhyay, S. Paul, Effect of target frequency, bias voltage and bias frequency on microstructure and mechanical properties of pulsed DC CFUBM sputtered TiN coating, Surf. Coat. Technol. 204 (2010) 3684e3697. [32] R. Gilmore, M.A. Baker, P.N. Gibson, W. Gissler, M. Stoiber, P. Losbichler, C. Mitterer, Low-friction TiNeMoS2 coatings produced by dc magnetron codeposition, Surf. Coat. Technol. 108e109 (1998) 345e351. [33] L.R. Shaginyan, M. Misina, J. Zemek, J. Musil, F. Regent, V.F. Britun, Composition, structure, microhardness and residual stress of WeTieN films deposited by reactive magnetron sputtering, Thin Solid Films 408 (2002) 136e147. [34] A. Nossa, A. Cavaleiro, Chemical and physical characterization of C(N)-doped WeS sputtered films, J. Mater. Res. 19 (2004) 2356e2365. [35] I. Martin-Litas, P. Vinatier, A. Levasseur, J.C. Dupin, D. Gonbeau, F. Weill, Characterisation of r.f. sputtered tungsten disulfide and oxysulfide thin films, Thin Solid Films 416 (2002) 1e9. [36] S.J. Bull, Failure modes in scratch adhesion testing, Surf. Coat. Technol. 50 (1991) 25e32. [37] M. Audronis, A. Leyland, P.J. Kelly, A. Matthews, Composition and structureproperty relationships of chromium-diboride/molybdenum-disulphide PVD nanocomposite hard coatings deposited by pulsed magnetron sputtering, Appl. Phys. A 91 (2008) 77e86. [38] J. Haider, M. Rahman, B. Corcoran, M.S.J. Hashmi, Deposition and characterization of hard-solid lubricant coating by closed-field magnetron sputtering, Surf. Coat. Technol. 200 (2005) 1080e1083. [39] B. Deng, Y. Tao, X. Zhu, H. Qin, The oxidation behaviour and tribological properties of Si-implanted TiN coating, Vacuum 99 (2014) 216e224. [40] P. Mutafov, M. Evaristo, A. Cavaleiro, T. Polcar, Structure, mechanical and tribological properties of self-lubricant WeSeN coatings, Surf. Coat. Technol. 261 (2015) 7e14.