Shear-Induced Structural Changes and Origin of Ultralow Friction of

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Shear-Induced Structural Changes and Origin of Ultralow Friction of Hydrogenated Diamond-like Carbon (DLC) in Dry Environment Praveena Manimunda,† Ala’ Al-Azizi,‡ Seong H. Kim,*,‡ and Richard R. Chromik*,† †

Department of Mining and Materials Engineering, McGill University, Montreal, QC H3A 0C5, Canada Department of Chemical Engineering and Materials Research Institute, Pennsylvania State University, University Park, Pennsylvania 16802, United States

‡

S Supporting Information *

ABSTRACT: The origins of run-in and ultralow friction states of a sliding contact of hydrogenated diamond-like carbon (H-DLC) and sapphire were studied with an in situ Raman tribometer as well as ex situ analyses of transmission electron microscopy (TEM), Raman spectroscopy, and nanoindentation. Prior to ultralow friction behavior, H-DLC exhibits a run-in period. During the run-in period in dry nitrogen atmosphere, the transfer film was formed and its uniformity and thickness as well as structure were varied. The duration and friction behaviors during the run-in depended on the initial surface state of the H-DLC coatings. A comparative study of pristine and thermally oxidized H-DLC revealed the role of surface oxide layer on run-in friction and transfer film formation. Attainment of the ultralow friction state appeared to correlate with the uniformity and structure of the transfer film evolved during the run-in, rather than its final thickness. TEM cross-section imaging of the wear track and the counter surfaces showed a trace of nanocrystalline graphite and a thin modified surface layer on both rubbing bodies. The comparison of hardness and reduced modulus of the wear tracks and the unworn surfaces as well as the ex situ Raman spectra suggested the densification of the wear track surfaces. Combining the in situ and ex situ analysis results, a comprehensive model was proposed for the formation and structure of the ultralow friction sliding contact of H-DLC. KEYWORDS: in situ Raman, ultralow friction, transfer film, hydrogenated diamond-like carbon

1. INTRODUCTION The superior mechanical and tribological properties of diamond-like carbon (DLC) coatings could be ideal in various engineering applications.1,2 Ranging from miniature mechanical devices to large scale automobile industry, DLC can be employed as a protective coating and a solid lubricant.3 The DLC coatings are classified according to the fraction of carbon−carbon sp3 bonds and hydrogen contents.2 Depending on application, properties of the DLC coatings can be tailored by doping metals (W, Ti, Cr, etc.) or non-metals (Si, O, N, F, B).4,5 Widely studied two variants of DLC are tetrahedral amorphous carbon (t-a-c) with ∼70% sp3 fraction and hydrogenated DLC (H-DLC) with 20−40 at. % hydrogen.6,7 The former is well-known for superior mechanical properties, and the latter is for ultralow friction in an inert atmosphere.8,9 The origin of the ultralow friction in H-DLC coatings under inert conditions is of great interest.9−13 It is noted that ultralow friction is always preceded by a high friction run-in period.9 It was thought that the formation of graphitic transfer film is the necessary criteria for ultralow friction, and such transfer films are produced during the run-in period.14,15 Alternatively, recent studies revealed the role of oxidized surface layers on the run-in period.9 A full understanding of the chemical and mechanical mechanisms responsible for ultralow friction has not yet been © XXXX American Chemical Society

achieved. This is in part due to the fact that most of the interpretations were based on ex situ characterizations.13,16,17 The chemical and structural changes spontaneously occurring at the DLC surface upon exposure to air complicate the data interpretation. 18 There are many reports of atomistic simulations or computational analysis of H-DLC, low friction sliding contacts.19−25 However, the available in situ spectroscopic data are very scarce for the H-DLC system. In situ Raman tribometry allows real time contact imaging and spectral analysis of the sliding contact.26−31 Tribological properties of most of the DLC coatings are often determined by third bodies.32 Recently, in situ Raman tribometry was used to study the contact dynamics of diamond-like nanocomposite coatings at ambient conditions and found the growth of thicker transfer film as a salient reason for the low friction steady state regime in amorphous diamond-like nanocomposite (DLN) coatings (C:H:Si:O) at ambient conditions.27 In situ Raman tribometry has shown that transfer films are important for many types of DLC coatings, but it has never been applied to H-DLC coatings that exhibit ultralow friction. Received: March 8, 2017 Accepted: May 1, 2017 Published: May 1, 2017 A

DOI: 10.1021/acsami.7b03360 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces In this paper, a custom-built in situ Raman tribometer was used to study the origin of ultralow friction in H-DLC samples at inert atmosphere. To gain more insight into the role of the oxidized surface layer and its structure in the run-in behavior, a comparative study was performed between pristine H-DLC and heat-treated H-DLC samples. The latter will be called thermally oxidized H-DLC (TO-H-TLC) hereafter. In addition to quantitative analysis of transfer film thickness, the chemical changes were identified using Raman spectroscopy. A series of ex situ analyses were also performed to understand the slidinginduced changes at the interface. Mechanical characterization and transmission electron microscopy (TEM) cross-section imaging of wear tracks were performed to support in situ contact analysis data and to gain deeper insight into the interfacial phenomena.

2. EXPERIMENTAL METHODS H-DLC coatings were produced on a silicon wafer substrate using plasma enhanced chemical vapor deposition (PE-CVD) at Argonne National Laboratory. The same coating was previously referred as NFC6.33 The deposition conditions are summarized in Table 1. In this

Figure 1. Schematic of the experimental configuration. Ball on a flat tribometer coupled with a Raman microscope. in the region of 1360 and 1560 cm−1 were fitted using a Gaussian function to extract the ID/IG ratio.36 The mechanical properties of the coating and the wear track were measured via nanoindentation (Triboindenter, Hysitron, USA). A diamond Berkovich indenter was used to indent the surface. The contact area of the Berkovich indenter was calibrated using standard fused silica sample. Indentation was performed at different normal loads (1−5 mN). The Oliver−Pharr method was used to fit the load− displacement curves.37 Line profiles of the wear tracks were measured using Wyko NT8000 3D optical interferometer (Veeco Instruments Inc., USA). TEM samples were prepared from the 300 °C TO-H-DLC wear track, sapphire hemisphere, and the pristine sample using a focused ion beam milling (Helios Nanolab 660, FEI USA) method. The cross section of the coating and wear track was imaged using a transmission electron microscope (Tecnai G2 F20 FEI, USA).

Table 1. H-DLC Deposition Conditions source gas composition (%) substrate

H2

CH4

RF power (W)

bias voltage (V)

gas press. (mTorr)

silicon

75

25

500−100

−500

10−13

study, pristine samples were the H-DLC coatings stored in ambient air. The TO-H-DLC samples were produced by heating H-DLC film at 180 and 300 °C in air for 2 h. The thermogravimetric analysis of the H-DLC samples were done using the TGA Q50 system (TA Instruments, New Castle, DE). A custom-built reciprocating ball on flat tribometer was used for the friction measurements.14 A sapphire hemisphere with a diameter of 6.35 mm was used as a counter surface. The H-DLC sample was mounted on to the reciprocating stage. A dead weight mechanism was used to apply normal load of 7.5 N, which gave an initial Hertzian contact pressure of 550 MPa. The track length and sliding speeds were 2.5 mm and 4 mm s−1, respectively. The tribometer was coupled with Leica DMLM free space microscope (Leica, Germany). A 20× ultralong working distance objective (SLMPLN/0.25) was used to image the contact. Using a video camera (Sony HDR-XR500 V), the contact region was recorded with a frame rate of 30 Hz. In order to maintain an inert atmosphere, the tribometer and the optical microscope were sealed with a bag, and dry N2 (ultrapure) was purged during the experiment. The snapshots were captured from the in situ contact image video frames. Using Newton’s rings method,34 the thickness of the transfer film was measured (Figure 1). The chemical analysis of the sliding contact was performed using a micro-Raman spectrometer (In Via, Renishaw, U.K). The Raman spectra were collected in back-reflection geometry using 20× objective lens with a numerical aperture of 0.25. A 514 nm argon ion laser was used as an excitation source. To avoid laser burning, neutral density filters were used to decrease the laser power to 5 mW. The laser spot size was ∼2 μm, and a 10 s data acquisition time was used to collect the spectra. The combination of 1800 lines/mm grating and 514 nm laser resulted in a spectral resolution ∼1 cm−1. During sliding test, the peak intensities and positions of the ID (∼1360 cm−1, A1g mode) and IG (∼1560 cm−1, E2g mode) bands were monitored.35 For in situ experiments, the spectral range was fixed at 800−2000 cm−1. The in situ Raman spectra recorded were an average of eight sliding cycles. For ex situ Raman analysis on the wear tracks and as-deposited coatings, the spectral range was fixed at 100−3200 cm−1. The spectra

3. RESULTS AND DISCUSSION 3.1. Characterization of Pristine and TO-H-DLC. The TGA data of H-DLC sample in air are shown in Figure 2a. With increase in temperature, weight gain was observed up to 300 °C. Above 300 °C sample weight started to decrease, and degradation of the sample was observed around 450 °C. Raman spectra recorded from the heat-treated samples are shown in Figures 2b and 2c. For the pristine H-DLC, an IG peak was observed at 1568 cm−1 and a broad shoulder was observed in the ID region (1355 cm−1). With increase in heat treatment temperature, an increase in ID band intensity and a blue-shift of the IG band were observed (Figure 2b). For 300 °C TO-HDLC, an ID peak was observed at 1358 cm−1 and IG at 1601 cm−1. Raman spectra of DLC coatings are sensitive to the sp2/sp3 ratio, clustering of sp2 phase, and bond length disorder.7,36,38,39 There are two characteristic peaks, known as G and D.40 The G band originates from the stretching of all pairs of sp2 atoms in both chains and rings.35 The breathing modes of sp2 atoms in ring structure give rise to the D band. Introduction of ring structures or local ordering within amorphous carbon matrix enhances ring structure breathing modes or D band intensity.41 The pristine H-DLC consists of amorphous carbon matrix and has weaker D band. The formation of the sp2-rich phase upon heat treatment has been reported.42,43 Growth of D band with heat treatment (Figure 2b) confirms formation of the sp2-rich phase. B


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Figure 2. (a) Thermogravimetric analysis of H-DLC sample in air (heating rate 5 °C/min). Raman spectra of pristine and heat treated H-DLC samples in the (b) D and G band region and (c) C−H stretch region. (d) Deconvolution of Raman spectra.

Figure 3. TEM cross-section image of (a) pristine H-DLC and (b) TO-H-DLC samples. In (b), the top right inset shows the high magnification image of the interface of the oxidized layer and the intact bulk, and the bottom right inset shows the EDS map of the TO-H-DLC sample.

Raman spectra recorded from the C−H stretching region were broad and asymmetric (Figure 2c). The decrease in C−H stretching mode intensity indicates loss of hydrogen upon heat treatment (Figure 2c) within the Raman probe depth. Since the C−H stretching bands are broad, it is difficult to deconvolute the contribution of CH-sp, CH2-sp2, or CH3-sp3. For 300 °C TO-H-DLC, an additional shoulder is observed at 3200 cm−1 (corresponding to second-order G).40 The deconvoluted Raman spectra in G and D region show the change in the spectral background due to photoluminescence background (PL) with heat treatment temperature (Figure 2d). The pristine samples showed large PL

background. The ratio of the slope of the spectral background (m) to the intensity of the G peak (Ig) can be used to estimate the hydrogen content:36 ⎧ ⎫ ⎪m ⎪ ⎨ [μm]⎬ H (at.%) = 21.7 + 16.6 log⎪ ⎪ ⎭ ⎩ Ig

(1)

Calculations using eq 1 revealed that the pristine sample had 36 at. % of hydrogen, whereas the 300 °C heat-treated sample showed 21% hydrogen. The 180 °C sample had intermediate hydrogen content (26 at. %). A study by Li et al. reported the loss of hydrogen during vacuum annealing of H-DLC.42 Their C


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Figure 4. In situ contact images recorded from sapphire−pristine H-DLC sliding contact. The growth of the transfer film is observed from changes in the Newton’s rings with increase in sliding cycles.

Figure 5. In situ contact images recorded from sapphire−TO-H-DLC sample. Loss of transfer film is observed during the run-in to superlubricious stage transition.

pristine and 300 °C TO-H-DLC samples as their hydrogen and oxygen content differed significantly. TEM cross-section images of pristine and TO-H-DLC samples are shown in Figures 3a and 3b, respectively. The coating thickness was about 1.2 μm. The 300 °C heat-treated sample showed an ∼185 nm thick surface layer with a slightly different contrast due to changes in atomic density upon oxidation.9 The top right inset in Figure 3b shows a finite width region of the interface between the oxidized layer and the intact bulk. The EDS map recorded from TO-H-DLC is shown in the bottom right inset of Figure 3b. A

findings showed slight reduction in hydrogen content up to 200 °C and large reduction above 300 °C.42 When thermal annealing was conducted in air, the H-DLC samples undergo thermally assisted relaxation as well as surface oxidation. The C−H bond or C−C bond dissociation is necessary to oxidize H-DLC. It is believed that both hydrogen release and thermal annealing are responsible for the bond rearrangement.41−53 In a previous study, XPS and NEXAFS analyses on thermally oxidized H-DLC samples showed increase in sp2 fraction and oxygen content (13.2%) for the 300 °C annealed sample.9 For detailed characterization and tribological studies, we choose D


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Figure 6. Variation of friction and transfer film thickness as a function of sliding cycles of sapphire ball on (a) pristine H-DLC and (b) TO-H-DLC surfaces.

Figure 7. In situ Raman spectra recorded for a sapphire ball sliding on (a) pristine H-DLC and (b) TO-H-DLC surfaces.

Figure 8. Variation of transfer film thickness, G band position (a, c), and ID/IG ratio (b, d) with sliding cycles for (a, b) pristine H-DLC and (c, d) TO-H-DLC.

uniform distribution of oxygen (red color) is seen in the top 185 nm of the TO-H-DLC sample. 3.2. Friction and in Situ Contact Analysis of Pristine and TO-H-DLC. Figures 4 and 5 show in situ contact images as

a function of sliding cycles recorded from the sapphire surface sliding against the pristine H-DLC and TO-H-DLC surfaces, respectively. In the case of pristine H-DLC, small patches of transfer film started to appear in the contact zone during 5−10 E


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Figure 9. SEM micrographs of (a) the sapphire contact region and (b) TO-H-DLC wear track. (c) Ex situ Raman spectra recorded from wear track and hemisphere surfaces.

cycles, and then the transfer film became thicker as the sliding continued. In the case of TO-H-DLC, a thicker transfer film was observed at the contact region during the first few cycles. The thick transfer film at the beginning of sliding is likely due to the oxidized surface layers which have a lower density and thus weaker than the bulk.9 From 4 to 20 cycles, the transfer film became thinner; at 60 cycles, a few patches were seen at the contact region. Above 100 cycles, the transfer film inside the contact region became uniform. Figure 6 plots the friction coefficient and transfer film thickness as a function of sliding cycle. In the case of pristine HDLC (Figure 6a), the transfer film thickness at the contact region increased to ∼100 nm in the first 10 cycles and then to ∼164 nm in the next 40 cycles. During this period, the friction dropped from 0.25 to 0.065. With further increase of sliding cycles, the transfer film thickness continued increasing gradually to 245 nm at 250 cycles, while the friction decreases marginally. Even at 658 cycles, the transfer film was not uniform yet (the last image of Figure 4), and the friction coefficient was still higher than 0.01. In the case of TO-H-DLC (Figure 6b), the friction coefficient was unstable and increased from 0.1 to 0.2 in the first 50 cycles and then dropped suddenly. During this period, the transfer film thickness dropped from 375 to 150 nm. Thermal oxidation resulted in the loss of hydrogen from the top surface region of H-DLC, and highly reactive intermediate sites formed during heat treatment might interact with the sapphire counter surface. The strong adhesive interaction between sapphire and the oxidized surface could lead to higher friction during the run-in period. Between 100 and 150 cycles, the coefficient of friction was ∼0.05, and some fluctuations were observed. The process of partial removal of the transfer film and filling of asperities could give momentary fluctuation in friction. Once the contact region attained uniform coverage of transfer film, the friction started to stabilize (from cycle no. 20 to 190 in Figure 5). Above 150 cycles, a sliding with ultralow friction (0.01) was observed. The ultralow friction and uniform transfer film (∼150 nm) was retained throughout the rest of the test duration (800 cycles). In situ Raman spectra recorded during the friction tests are shown in Figure 7. Figure 8 compares the transfer film thickness and changes in the G-band position as well as ID/IG ratio of the Raman spectra as a function of sliding cycle. The pristine H-DLC contact showed a sudden increase in the G band position and the ID/IG ratio in the first 50 cycles, followed

by a gradual and continuous increase in the subsequent sliding cycles (Figures 8a and 8b). These changes in Raman spectral features with sliding cycles appeared to correlate with the trend in the transfer film thickness change for the sapphire−pristine H-DLC contact. However, such a correlation is not observed for the sapphire−TO-H-DLC contact (Figures 8c and 8d). The G band position of the film in the sapphire−TO-H-DLC contact dropped suddenly to ∼1560 cm−1 in the fist few cycles and then increased back to ∼1590 cm−1 at ∼10 cycles. During the subsequent ∼100 cycles, it remained unchanged while the transfer film thickness decreased by ∼50%. Regardless of the transfer film thickness change, the common features of the in situ Raman spectra for both pristine and TO-H-DLC films were that the G band position increased to 1590−1600 cm−1 and the ID/IG ratio asymptotically approached 0.75 ± 0.02 as the sliding cycle continues even though their initial values before sliding were different. The shifting of G band and Id/Ig ratio indicated the sp2-rich transfer film. Previously, it was believed that graphitization of the thicker transfer film is necessary to achieve ultralow friction in the DLC system.29,17 However, the in situ contact images (Figures 4−6) and in situ Raman analyses (Figures 7 and 8) suggest that the key determinant of the superlubricious state is not the transfer film thickness, but the type and uniformity of the transfer film. At a first glance, the shift of the “G” band position toward 1600 cm−1 and the higher ID/IG ratio could be interpreted as the possibility of nanocrystalline graphite formation.39,40 However, it was difficult to predict whether the transfer film has undergone the full transformation from the in situ Raman study alone. 3.3. Ex Situ Analysis of Wear Tracks and Transfer Films of TO-H-DLC. Since the sapphire−TO-H-DLC contact reached the ultralow friction in our test condition, their wear track and transfer films were further analyzed ex situ. SEM micrographs of the sapphire counter surface and TO-H-DLC wear track are shown in Figures 9a and 9b. Thin patches of transfer film were observed in the contact region (marked by a dotted circle). Debris were seen at the edges of the wear track. The TO-H-DLC wear track was ∼180 nm deep (see Supporting Information for profilometry data). Raman spectra recorded from the TO-H-DLC wear track and the transfer film on the counter surface are compared in Figure 9c. The transfer film Raman spectra showed a stronger D band and a larger shift in the G band position (1602 cm−1) compared to the original TO-H-DLC surface before sliding (Figure 2). The IG band F


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Figure 10. TEM cross-section images from TO-H-DLC wear track and sapphire hemisphere contact region. (a) TO-H-DLC wear track center. (b) Edge of the wear track where debris or transfer films were found. (c) Higher magnification image showing the presence of nanocrystalline graphite at the transfer film-TO- HDLC interface. (d) Diffraction pattern from the region where nanocrystalline graphite was spotted. (e) Cross-section TEM micrograph of sapphire hemisphere. (f) Higher magnification image showing structural differences across the transfer film on sapphire surface. A denser compacted layer was observed at the interface.

position (1602 cm−1) and the higher ID/IG ratio (0.76) indicated the presence of more ordered sp2 clusters within the transfer film on the sapphire counter surface. The wear track region had a lower ID/IG ratio (0.62), indicating higher sp3 content. Full range spectra recorded from wear track and debris region are given in the Supporting Information. The Raman spectra of the TO-H-DLC wear track had higher PL background (see Supporting Information), indicating a higher hydrogen content. The TEM cross-section images recorded from the TO-HDLC wear track and hemisphere surface are shown in Figure 10. Focused ion beam milling was used to cut out two crosssectional slices: one at the center of the wear track and the other at the wear track end where debris was found (see

Supporting Information for SEM micrographs). The wear track region (Figure 10a) had about 90 nm thick modified layer after the friction test for 800 cycles, whereas the debris region (Figure 10b) had much thicker oxidized layers onto which a thick layer of transfer film containing traces of graphitic flakes was deposited. The high-resolution image (Figure 10c) and electron diffraction patterns (Figure 11d) suggested that they are nanocrystalline graphite. Since these nanocrystalline graphite flakes were seen at the interface of transfer film and the TO-H-DLC surface, one could infer that nanocrystalline graphite might have formed during the early sliding cycles when thicker transfer film was present at the sliding interface. However, it is important to note that most contact regions (wear track) did not show such nanocrystalline graphite flakes. G


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Figure 11. (a) Comparison of hardness and reduced modulus as a function of indentation depth in the unworn and wear track regions of TO-HDLC. (b) Representative indentation load vs displacement curves.

n2 = −0.13 + 3.7 × 10−2FWHM G [cm−1]

Figure 10e shows the TEM cross-section image of sapphire contact region. The thickness of the transfer film varied from 25 to 50 nm at different locations of the sapphire contact region. It is noted that the sapphire asperities were fully filled with HDLC transfer film. The high magnification image (Figure 10f) shows two distinct zones within the transfer film on the sapphire surface. An ∼10 nm thick compact region is seen at the sapphire transfer film interface, followed by a loose transfer layer on the top. The mechanical properties of the TO-H-DLC wear track were measured using nanoindentation and then compared with those of the unworn TO-H-DLC surface (Figure 11). For the unworn region, the hardness and reduced modulus values were very low at a small indentation depth and increased toward the values of the pristine H-DLC surface (H = 7.2 GPa, Er = 75 GPa) with the increase of indentation depth. This implied that the top surface layer was affected by thermal oxidation and the bulk properties were unchanged. This is different from the previous study reporting a decease in hardness of the entire sample upon annealing in a vacuum.54 In the case of wear track, nanoindentation measurements showed hardness and reduced modulus higher than the pristine H-DLC values up to 100 nm of indentation depth. When the indentation depth was larger than 100 nm, the hardness and reduced modulus started to drop toward the values of the pristine sample. The increase in hardness observed in the TOH-DLC wear track may suggest a hypothesis of top layer densification. Raman spectra recorded from the TO-H-DLC wear track showed higher sp3 content. The in situ contact images (Figure 5) showed change in color after 60 cycles. The color change indicates a refractive index change, whereas higher hardness is indicative of surface densification. Both phenomena must be associated with structural modifications. To test this hypothesis, ex situ Raman spectra were processed with the model suggested by Casiraghi et al. relating the FWHM of the G peak to the density and refractive index:36 ⎛ g ⎞ ρ ⎜ 3 ⎟ = 0.257 + 0.011 × FWHM G [cm−1] ⎝ cm ⎠

(3)

Using the G band fwhm of the spectra shown in Figures 3 and 9, the density of the unworn and wear track regions was estimated to be ∼1.44 and ∼1.53, respectively. The refractive index estimated from the Raman spectra for TO-HDLC surface and wear track are 1.74 and 1.87, respectively. The refracvtive index of the H-DLC samples are sensitive to sp3/sp2 fraction. High density and refractive index confirm structural modifications (increase in sp3 content) on the wear track surface. However, it is not clear whether the changes took place during sliding or heat treatment process. Molecular dynamic simulations showed that shear-induced strain localization can induce rehybridization of sp2 species to sp3 in amorphous carbon system, which could result in increase in density and improved mechanical properties.55 A study by Li et al. showed vacuum annealing-induced structural changes in H-DLC samples.54 When annealing was done in air, surface oxidation takes place additionally. During oxidation, H-DLC sample loses hydrogen atoms from the surface. However, the growth of thicker oxide layer could prevent further loss of hydrogen from the subsurface. In TO-H-DLC sample, the oxygen intake was limited to top ∼185 nm (see Figure 3b, EDS map), and highresolution TEM cross-section images showed a well-defined interface (Figure 3b, inset). It is speculated that at this interface the released hydrogen (from subsurface) reacts with sp2 carbon to form hydrogenated sp3 sites. The possible reactions during hydrogen interaction with carbon sp2 sites are hydrogenation or hydrogen abstraction.56−58 Molecular dynamics study by Jariwala et al. showed interaction of hydrogen with sp2 and sp3 carbon in hydrogenated amorphous carbon.59 They used the Eley−Rideal type mechanism to explain the hydrogenation of sp2 sites.59 Bonding of hydrogen with sp2 carbon atom having configuration similar to the one in graphene changed hybridization to sp3.59 Hydrogenation changed bond energies of neighboring sp2 carbon atoms from −5.03 to −3.69 eV.59 Ex situ analysis showed the higher sp3 and hydrogen content in the wear track of TO-HDLC (Figure 9c). Since the hydrogenation reaction was exothermic in nature59 with lower activation barrier, it is

(2) H


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Figure 12. Schematic illustration of run-in and steady-state ultralow friction sliding contact of the H-DLC film (not to scale). Transition from the run-in to ultralow friction sliding involves removal of the oxidized layer, asperity filling, ordering of sp2 clusters within the transfer film, and compaction of wear track surface. Hydrogen-terminated sp3 sites on wear track provide surface passivation.13

H-DLC slides against sapphire, the intermediate reactive sites created during the surface oxidation can interact strongly with the counter surface and form well-adhered transfer film. As seen in Figure 10e, the sapphire asperities are filled with the transferred film. The majority of the surface oxide layer was removed during the run-in (Figures 5 and 10). Even though, oxidized layer was hydrogen deficient, removal of surface oxide layer exposes hydrogenated sp3-hybridized carbon cites. Ex situ Raman analysis confirmed hydrogen-rich wear track (Figure 9c). Even though the starting sliding pair is not DLC vs HDLC, the formation of uniform compact transfer film transforms the sliding interface into H-DLC vs H-DLC type contact. Once we achieve compacted transfer film at the interface, the surface passivation mechanism proposed by Erdemir appears to pertain to the H-DLC surface.13,60 It is believed that hydrogen passivates the surface dangling bonds and lowers adhesive interaction between sliding surfaces.13 A study by Erdemir et al. demonstrated the role of hydrogen in lowering the friction of non-hydrogenated DLC sliding contact.12 Their time-of-flight secondary ion mass spectrometry (ToF-SIMS) study confirmed the formation of hydrogen passivation layer on wear track.12 A recent study by Wang et al. explained the role of adsorbed nitrogen on low friction sliding of hydrogenated amorphous carbon.62 However, in our study we did not find any evidence to support the role of adsorbed nitrogen on low friction. Even though many researchers attempted to solve the low friction characteristics of H-DLC, the long pending question in H-DLC sliding was actual sliding interface. In situ and ex situ analysis presented in this study shed light into the H-DLC sliding interface.

fair to assume the possibility of such reaction in the TO-HDLC system. Along with interface chemistry, one must understand the contact dynamics to achieve full understanding of ultralow friction sliding contact. In situ contact imaging revealed contact dynamics and transfer film formation (Figures 4 and 5). Even though H-DLC samples are known for their ultralow friction sliding in inert atmosphere, in our experiments they showed slightly higher friction (∼0.05) than TO-H-DLC samples. In an earlier study, ultralow friction was reported for H-DLC−steel sliding contact under inert conditions.9 The observed difference in friction characteristics was related to the velocity accommodation modes. The velocity accommodation mode associated with ultralow friction sliding of TO-H-DLC was interfacial sliding between compacted transfer film and wear track. During sapphire−pristine H-DLC sliding we did not observe uniform, compacted transfer film at the interface. Instead, patches of transfer film were seen at the interface, and their thickness varied during steady state friction sliding (Figures 4 and 6a). In the absence of compacted transfer film, it is possible to have both transfer film shearing and interfacial shear. The dominance of transfer film shearing may lead to larger energy dissipation at the sliding interface and thus higher friction. Formation of compacted transfer film depends on experimental conditions. A recent study by Liu et al. showed the influence of tribofilm on low friction sliding of hydrogenated amorphous carbon.61 Their findings also showed loose transfer film on steel surface at lower loads and compacted transfer film above 124 MPa (which showed lower friction).61 Combining all in situ and ex situ analysis results of this study, a schematic illustration of the sliding contact during run-in and superlubricious state is proposed in Figure 12. When a thermally oxidized I


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4. CONCLUSIONS The combination of in situ analyses of the transfer film formation during the run-in period and ex situ analyses of the transfer film responsible for the ultralow friction state revealed the concerted effects of multiple dynamic changes occurring in the sliding contact of H-DLC in dry nitrogen. During the runin, the H-DLC surface wears and the transfer film is formed on the counter surface. The thickness and uniformity of the transfer film during the run-in as well as the friction and duration of the run-in varied depending on the initial oxidation state of the H-DLC surface. Oxide layer formed during annealing in air appeared to prevent hydrogen loss from the subsurface. Annealing-induced sp2 clusters reacted with released hydrogen to form sp3 sites at the oxide−H-DLC interface. Formation of sp3 sites increased local density. Nanoindentation performed on TO-H-DLC wear track showed increase in hardness. Cross-section TEM images of the mating surfaces confirmed surface densification. Interfacial sliding between compacted transfer film and dense hydrogenated wear track produced ultralow friction.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b03360. Detailed SEM, TEM micrographs, 3D optical profilometry image and Raman spectra recorded from TO-HDLC wear track and debris (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected]; Tel +1 514-398-5686. *E-mail [email protected]; Tel +1 814-863-4809. ORCID

Praveena Manimunda: 0000-0002-1851-4777 Seong H. Kim: 0000-0002-8575-7269 Richard R. Chromik: 0000-0003-4052-1858 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS P.M. acknowledges technical assistance from Line Mongeon, Wu, Ken and David Liu of the Facility for Electron Microscopy Research at McGill University. P.M. and R.R.C. thank the Natural Science and Engineering Research Council (NSERC) of Canada Discovery Grants program for financial support. A.A. and S.H.K. were supported by the National Science Foundation (Grant No. CMMI-1131128 and 1435766). The authors acknowledge Osman Eryilmaz and Ali Erdemir at Argonne National Laboratory for kindly providing DLC samples used in this study.

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REFERENCES

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Shear-Induced Structural Changes and Origin of Ultralow Friction of

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