Understanding the run-in tribological behavior of amorphous carbon films against Al₂O₃ balls under high vacuum

Abstract

Run-in behavior of amorphous carbon films significantly affected their tribological properties. Generally, for amorphous carbon films, a low friction coefficient was obtained after a run-in period during which the friction coefficient was high and gradually decreased to a low value under vacuum. However, in this paper, the friction coefficient was initially low and gradually increased to a high value during run-in period under vacuum. According to experimental results, first principles calculations were selected to probe the possible mechanism of the anomalous run-in behavior. The results showed that the weak and strong atomic interactions should be attributed to low and high friction.

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1. Introduction

Recently, studies suggested that amorphous carbon films exhibited low friction in several environments, such as high vacuum, dry N₂ and ambient air.^1–3 Causing low friction in the intermediate friction stage is mainly attributed to the formation of a transfer film. Hirvonen and Liu declared that low friction of amorphous carbon films in the steady-state depended on the formation of a carbon-rich transfer film on the counterface.^4,5 Although the frictional properties and the wear mechanism of amorphous carbon films have been investigated by many researchers, the tribological behavior of amorphous carbon films sliding against different counterparts has yet to be well understood so far.

As we known, the different counterparts have an effect on the tribological properties in vacuum. In our previous work,⁶ the amorphous carbon film sliding against Al₂O₃ ball exhibited low friction, while sliding against Si₃N₄ ball displayed high friction under high vacuum. The other interesting result was that the amorphous carbon film sliding against Si₃N₄ ball exhibited super-low friction (<0.01), while sliding against Al₂O₃ ball displayed high friction (>0.1) at run-in processes. It was often neglected much attention to these phenomena, like the other researches.^7–9 In general, the run-in period was a transient state where the system adjusts before reaching steady state. Following the run-in period, the system achieved a steady state where the friction coefficient and wear rate tended to level off.¹⁰ Actually, no wear for self-mated hydrogenated amorphous carbon film was observed sliding in the n-propanol-containing humid air during run-in period. And the run-in behaviors were significantly influenced by environments as stated by Marino.¹¹ Besides, the run-in behaviors were governed by the steady transfer film, and the tribological systems could fast exhibit low friction due to the formation of steady transfer film.^12,13 However, the run-in behaviors of the amorphous carbon film under vacuum were not considered in these researches. Therefore, a detailed study for run-in behaviors of amorphous carbon film under vacuum would have deeply understood the friction and wear behaviors.

There existed the difficulties in direct observation of the atomic interaction processes by in situ experiments with atomic resolution. However, first principles calculations provided a powerful tool to capture atomic details and gain a deeper insight into the run-in tribological behavior of amorphous carbon film at the nanoscale. Especially, the bonding patterns in the carbon films could affect the tribological performances.¹⁴ Gueorguiev and coworkers theoretically predicted and guided the synthesis of carbon films based on first principles calculations.^15,16 And they pointed out the bonding patterns in the carbon films, which were conducive to probing the run-in behavior of carbon films. To our knowledge, the amorphous carbon film surface was often represented by a diamond surface following the common practice used in the literature of employing diamond to use as a model to study the amorphous carbon surfaces.^17,18 Since Righi and coworkers reported the adhesion and shear strength of diamond (001) surfaces based on first principles calculations.^19–21 Diamond (001) surface had attracted more attention due to a stable reconstruction constituted of dimers based on first principles calculations. And they believe that the value obtained for the work of adhesion will be more realistic if the surface reconstruction is taken into account. Based on their investigations, we carried out the calculations on the O-/Al-terminated Al₂O₃(012)/diamond(001) interfaces based on first principles calculations in order to probe the possible run-in mechanism. Understanding the origin of the run-in tribological behaviors of the amorphous carbon films was important to design optimum operation conditions to prevent potential failures of amorphous carbon films.

The paper is organized as follows. Materials and methods are illustrated in section 2. In section 3.1, we describe the tribological behavior of the amorphous carbon film against Al₂O₃ ball. In section 3.2, we illustrate the results of the characterization of the wear debris using SEM, Raman and XPS. In section 3.3, we describe the results of the first-principles calculations. In section 4, we discuss the possible mechanism of run-in behavior. The main conclusions are listed in section 5.

2. Materials and methods

2.1 Experimental section

The amorphous carbon film was deposited on polished stainless steel substrates in a d.c.-pulse parallel-plate hollow-cathode system. Fig. 1a showed the deposition process of the amorphous carbon film. To enhance the adhesion of the amorphous carbon film, we first coated the substrates with a silicon interlayer. A Si interlayer of about 200 nm was deposited with SiH₄ gas of 50 sccm and Ar gas of 100 sccm at 13.0 Pa. An amorphous carbon film was deposited in a SiH₄, CF₄, C₂H₂ and Ar environment. The F_x1–Si_y1–DLC layers were deposited in SiH₄ (25 sccm), CF₄ (25 sccm), C₂H₂ (150 sccm) and Ar (100 sccm) at 4.0 Pa. The F_x2–Si_y2–DLC layers were deposited in SiH₄ (25 sccm), CF₄ (25 sccm), C₂H₂ (100 sccm) and Ar (100 sccm) at 2.8 Pa. Finally, a pure DLC layer was deposited on the F–Si–DLC layers surface. The pure DLC layer was deposited from C₂H₂ (150 sccm) and Ar (100 sccm) gases by the same deposition system. The substrate bias voltage was maintained at −800 V, a duty cycle of 30%, and a repetition frequency at 1.5 kHz. No external heating of the substrate was employed, and the maximum temperature during deposition was about 180 °C. As shown in Fig. 1b, an amorphous carbon film of 10 μm in thickness is deposited on a Si interlayer of 200 nm in thickness.

Fig. 1 (a) The deposition processes of the amorphous carbon film. (b) The SEM cross-sectional image of the amorphous carbon film.

The fractured cross-sectional morphology, EDS and the thickness of the as-deposited film were examined using a JSM-6701F cold field scanning electron microscope (SEM). The chemical composition of debris around the wear scar of Al₂O₃ ball was examined using X-ray photoelectron spectroscopy (XPS : AXIS ULTRA ^DLD). The measurement was carried out using 36 MeV iodine ions as incoming ion projectile. The diameter of the beam spot was 15 μm. Raman spectra of the amorphous carbon film was obtained by a Horiba Jobin Yvon LABRAM-HR800 spectrometer using an excitation wavelength of 532 nm. The typical spectrum was recorded in the range of 800–2000 cm⁻¹, data acquisition time was 60 s and diameter of beam spot was 0.2 μm. The contact 2D surface profiler image of the wear track was performed using AlphaStep® D-100 Stylus Profiler. Besides, the friction tests were performed using a ball-on-disk apparatus under high vacuum (2 × 10⁻⁴ Pa). The tested samples were amorphous carbon film deposited on polished stainless steel substrates. The Al₂O₃ and Si₃N₄ ball of 4 mm in diameter with Ra ≤ 14 nm was used as the counterpart. All the tests were done at a sliding speed of 300 rev min⁻¹ corresponding to linear speed of 0.125 m s⁻¹, a rotational radius of 6 mm, and under a normal load of 2 N, corresponding to a initial theoretical Hertzian contact pressure of 0.63 GPa.

2.2 Computational section

The first principles calculations were carried out using CASTEP code,²² which used the energy plane-wave pseudopotential total calculation method based on the density functional theory. The generalized-gradient approximation with the Perdew–Burke–Ernzerhof exchange–correlation functional was selected for all the calculations. A plane-wave cutoff energy of 350 eV and Monkhorst–Pack k-point meshes with a density of (7 × 7 × 1) were employed throughout. The energy tolerance was 2.0 × 10⁻⁵ eV per atom, the force tolerance was 0.05 eV Å⁻¹, the displacement tolerance was 0.002 Å and maximum stress was 0.02 GPa. Firstly, the calculated results were obtained at 0 K. The diamond (001)-(2 × 2) surfaces and O-/Al-terminated Al₂O₃ (012) surfaces were relaxed to a minimum energy configuration, as shown in Fig. 2. During first principles calculations, the diamond (001)-(2 × 2) surface did not reconstruct. However, interestingly, during the calculation of the O-/Al-terminated Al₂O₃(012)/diamond(001) interfaces, the diamond (001)-(2 × 2) surfaces reconstructed, which indicated that the value obtained for the work of adhesion would be more reliable according to ref. 19–21. Then, to minimize the mismatch between constituent slabs and meanwhile achieve smallest interface areas, specific crystal forms and crystal faces were chosen to construct interfaces leading to the mismatch less than 3.0%, which were illustrated in Fig. 3. The super-cell (2 × 2) was used in all calculations. Then the work of separation (W_sep) could be computed viawhere the total energy of system (E_tot) was calculated by letting the atoms relax in their initial positions without allowing the constrained interface structure and atoms to relax. The change of E_tot corresponded to be relative to that of reference state E_tot′ at the far separated interface of 9.0 Å. The relative energy change ΔE_tot = E_tot − E_tot′ represented the diamond/Al₂O₃ interfaces. A represented the interface area. The top slab was moved to various positions by a lateral spatial increment of 0.5 Å along the sliding direction, while the bottom slab was fixed. Following each increment of the upper slab, the interface system was relaxed to a minimum energy configuration. Periodic boundary conditions were applied in the x–y plane and a large vacuum thickness of 20 Å was added along the z direction. The sliding direction was set along the x axis. The E_tot′ remained a constant during sliding. We had τ = (∂σ_int/∂x)|_max = (−∂W_sep/∂x)|_max, τ was the shear stress along sliding direction.²³

Fig. 2 Diamond (001)-(2 × 2) surfaces: (a) original configuration and (b) relaxed configuration. The parts marked as a black dashed box are fixed to represented bulks of the slabs and the rest parts are allowed to relax. After relaxation, the diamond (001)-(2 × 2) surface does not reconstruct. O-terminated Al₂O₃ (012)-(2 × 2) surfaces: (c) original configuration and (d) relaxed configuration. Al-terminated Al₂O₃ (012)-(2 × 2) surfaces: (e) original configuration and (f) relaxed configuration.

Fig. 3 Interface structures of (a) Al-terminated Al₂O₃(012)/diamond(001), (b) O-terminated Al₂O₃(012)/diamond(001). Space group for Al₂O₃ is R3̄c. The parts marked as a black dashed box are fixed to represented bulks of the slabs and the rest parts are allowed to relax. We focus on the part of “area C” during the relaxed processes.

3. Results

3.1 Tribological behavior

Fig. 4 shows the friction coefficient curves of amorphous carbon film against Al₂O₃ ball. On the basis of the characteristic of the curve, we divide the run-in period into two stages marked as “A” and “B” in Fig. 4. The friction coefficient initially decreases, while increases closely to 0.3 after 700 sliding cycles, then decreases after the stage B, which shows an unsteady tribological behavior rather than keeping a steady state value over time. The similar unsteady behaviors are observed unpredictably in repetition tests (inset image in Fig. 4).

Fig. 4 Friction coefficient curves of the amorphous carbon film sliding against Al₂O₃ ball in high vacuum (2 × 10⁻⁴ Pa), inset is the repeated friction test.

Fig. 5 gives the wear depth of the amorphous carbon film sliding against Al₂O₃ ball in high vacuum (2 × 10⁻⁴ Pa). The wear depth is merely about 250 nm, which indicates that the amorphous carbon film reveals superior tribological properties against Al₂O₃ ball under high vacuum. The similar wear resistance property is observed unpredictably in repetition tests (inset image in Fig. 5).

Fig. 5 The wear depth of the amorphous carbon film sliding against Al₂O₃ ball in high vacuum (2 × 10⁻⁴ Pa), inset is the repeated friction test.

For comparison, Fig. 6a represents the friction coefficient of the amorphous carbon film against Si₃N₄ ball in high vacuum (2 × 10⁻⁴ Pa). The friction coefficient initially remains a low value of less than 0.05 until 400 sliding cycles and then increases steadily with increasing sliding cycles. Finally, the film is worn out. When the sliding cycles is 250 and the friction coefficient is more than 0.2, we stop the friction tests. Then the worn surfaces are analyzed by SEM, as shown in inset of Fig. 6a and b. As shown in inset of Fig. 6a (low friction region), less carbon debris are observed on the worn surface. Fig. 6b (high friction region) shows that the large amount of carbon debris is found on/around the worn surface.

Fig. 6 (a) Friction coefficient curves of the amorphous carbon film sliding against Si₃N₄ ball in high vacuum (2 × 10⁻⁴ Pa), inset is the SEM image of worn surface of Si₃N₄ ball (black dashed line marked region). (b) SEM image of worn surface of Si₃N₄ ball (red dashed line marked region), insets are enlarged images.

3.2 SEM, Raman and XPS analysis

In order to probe the possible run-in behavior, we have to characterize the wear debris around the ball surface. Fig. 7a shows a SEM examination of the Al₂O₃ ball surface that contacts with the amorphous carbon film. Fig. 7b–f shows EDS elemental maps taken from the Al₂O₃ ball surface, which indicate that the wear scar mainly consists of Al and O. It reveals that the carbon-rich transfer film does not form on the wear scar.

Fig. 7 Secondary electron (SEI) of the wear scar of Al₂O₃ ball for the first stage (“A”) in Fig. 4. The elemental EDS maps are taken from the whole areas shown in (a) are shown for (b) Al K, (c) O K, (d) C K, (e) F K and (f) Si K.

For the second stage (“B”), a SEM examination of the Al₂O₃ ball surface (Fig. 8a) shows that the wear scar also mainly consists of Al and O. Differently, the wear debris consisted of Al, O, C, F and Si (Fig. 8b–f) is detected at the circumference of the wear scar. Compared to Raman spectra of original film, Raman spectra (Fig. 9), which take from the wear scar (Fig. 7a), show that the signal of carbon debris is very weak. The signals of carbon on the wear scar shown in Fig. 8a are not detected by Raman spectra.

Fig. 8 Secondary electron (SEI) of the wear scar of Al₂O₃ ball after the second stage (“B”) in Fig. 4. The elemental EDS maps are taken from the whole areas shown in (a) are shown for (b) Al K, (c) O K, (d) C K, (e) F K and (f) Si K.

Fig. 9 Raman spectra of original film and wear scar acquired in the area in Fig. 7a.

Subsequently, the chemical composition of the wear debris shown in Fig. 8 is analyzed by XPS to probe the possible tribochemical reactions between two contact faces. Fig. 10 shows the XPS C 1s spectrum of the wear debris at the circumference of the wear scar. Fitting C 1s spectrum of the amorphous carbon film allows discrimination of these four components at the binding energies 283.7 eV, 285.2 eV, 286.1 eV, 287.4 eV corresponding to Si–C bonds, C–C bonds, C–O bonds and C=O bonds, respectively.^24–27 Fig. 11 shows the fitting XPS Al 2p spectrum taken from the same wear debris. The binding energies of 72.4 and 74.9 eV are assigned to Al₂O₃.^28,29

Fig. 10 The fitting XPS C 1s spectrum of the wear debris around the wear scar shown in Fig. 8.

Fig. 11 The fitting XPS Al 2p spectrum of the wear debris around the wear scar shown in Fig. 8.

3.3 First-principles calculations

Fig. 12 shows the interfacial structures for O-terminated Al₂O₃(012)/diamond(001) interface during sliding. The bond style of C and O transforms from single to double during sliding. And the O atom transfers to the diamond surface when the sliding distance is 1.5 Å. As reported by Sanderson,³⁰ the bond energy of the double bond is as much as 1.488 times than that of a single bond. Thus it would dissipate more energy to break the C=O bonds during sliding. The W_sep values are shown in Fig. 12 as well. The large W_sep value for O-terminated Al₂O₃(012)/diamond(001) interface during sliding results from the strong C–O or C=O bonds. The spatial variation of ∂σ_int/∂x with sliding distance is shown in Fig. 13, and a linear interpolation results of ∂σ_int/∂x and W_sep are also shown in Fig. 13. (∂σ_int/∂x)|_max and W_sep^max are about 0.68 GPa and 1.56 J m⁻² for Al-terminated Al₂O₃(012)/diamond(001) interface during sliding, respectively.

Fig. 12 Interfacial structures for O-terminated Al₂O₃(012)/diamond(001) interface during sliding. W_sep values in J m⁻² are shown in the parentheses below the respective structures.

Fig. 13 W_sep and variation of ∂σ_int/∂x along the sliding direction for Al-terminated Al₂O₃(012)/diamond(001) interface.

4. Discussion

The amorphous carbon film is successfully fabricated using a d.c.-pulse parallel-plate hollow-cathode system. The film exhibits a variable run-in tribological behavior under high vacuum. In the first stage (“A”), the film displays low friction, while the film exhibits high friction (>0.1) under the second stage (“B”). Differently, for hydrogenated amorphous carbon film, a friction coefficient is obtained after an initial induction period (called a run-in period) during which the friction coefficient is high and gradually decreases to a low value as the sliding cycles repeats, as reported by Erdemir.³¹ In general, the amorphous carbon film surface wears off and carbon transfer films are often formed on the counterpart surface during the run-in period.^32–35 However, the carbon-rich transfer film would lead to high friction and wear under vacuum, as declared by the most of previous studies.^1,6 In this paper, for the first stage (“A”), SEM and Raman spectra shows that there is several carbon wear debris on the wear scar as shown in Fig. 7. From Fig. 6, the carbon debris results in high friction under high vacuum. In deed, Yoon et al. declared that the quantity of carbon wear debris was closely related to the high friction under high vacuum. The wear debris was much less observed on the worn surface, the low friction coefficient obtained.³⁶ Thus, the low friction coefficient (“A”) may be attributed to a spot of carbon debris on the worn surface shown in Fig. 7. The XPS analysis obtained from Fig. 8 reveals that the wear debris is mainly composed of C and Al₂O₃. Thus, the anomalous run-in behavior is possibly related to these wear debris. To our knowledge, incorporation of Al₂O₃ particles into epoxy and aluminum alloy, the friction coefficient was often more than 0.3, as reported by the most previous studies.^37–39 And Al₂O₃ wear particles produced during the friction test also could lead to high friction coefficient.⁴⁰ Importantly, the Al₂O₃ sliding against Al₂O₃ exhibited high friction coefficient under vacuum.⁴¹ Thus, the high friction coefficient at the second stage (“B”) may be attributed to the amorphous carbon and Al₂O₃ wear particles. However, how to understand the run-in behavior at the atomic scale, which is some difficulty in directly observing from the sliding experiments. Thus, the first-principles method is selected to probe the possibly run-in mechanism. Of course, although the first-principles calculations and sliding experiments do not performed under identical contact and environmental conditions, they could depict a possible run-in mechanism through complement each other. From Fig. 12, the W_sep ranges from the values of 6.77 J m⁻² to 7.68 J m⁻². Such high W_sep value resulted from the instability of the O-terminated Al₂O₃ surface.²³ Qi reported that high W_sep value corresponded to high adhesion, and high adhesion between two contact faces led to high friction during sliding.¹⁷ Thus, in Fig. 4, the high friction occurring in the second stage (“B”) may result from the high adhesion. As shown in Fig. 13, the low friction should be attributed to low W_sep, τ and weak interactions between C and Al atoms.^17,42

5. Conclusion

In this work, the film exhibits variable run-in tribological behavior under high vacuum. The first principles calculations are used to probe the possible mechanism of the run-in behavior. The calculated results show that the variable run-in behavior mainly depends on the atomic interactions between two contact faces. The low friction is attributed to the weak interactions between C and Al atoms resulting in low adhesion. The high friction depends on the strong interaction between C and O atoms leading to strong adhesion.

Acknowledgements

The work was supported by the National Nature Science Foundation of China (Grant 51322508 and 11172300) and Nature Science Foundation of Gansu Province of China (Grant 145RJDA329). The authors gratefully acknowledged Mr. Yaonan Zhang and Guohui Zhao for supporting the numerical simulation at Lanzhou branch of supercomputing CAS.

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