1 Introduction
Molybdenum disulfide (MoS
2) solid lubricants have historically been used as coatings in space applications because of their low friction coefficients (
µ < 0.05) and wear rates (
k< 1 × 10
–6 mm
3/Nm) in inert and vacuum environments [
1‐
3]. Pure MoS
2 coatings are not commonly used in terrestrial applications due to high wear rates and oxidation when exposed to water and oxygen [
4‐
11]. To mitigate adverse interactions in terrestrial environments, dopants such as C, Sb
2O
3, Au, Ni, Ta, and Ti are commonly added to improve the tribological performance and environmental robustness [
12‐
20]. While reported mechanisms for performance improvement of composite coatings vary depending on additives, composites such as MoS
2/Sb
2O
3/Au have been shown to help facilitate the expression and retention of MoS
2 at the interface through agglomeration of Au nanoparticles [
21]. Other composites such as MoS
2/C/Sb
2O
3 have been shown to exhibit extremely low shear strengths. The sliding interface of MoS
2/C/Sb
2O
3 changes during sliding depending on the environment, with the surface becoming carbon-rich in humid testing conditions and MoS
2-rich in dry, inert environments [
22].
A common hypothesis for the improved environmental
resilience of composite MoS
2 coatings involves densification and hardening imparted by dopants [
23‐
28]. Sputtered pure MoS
2 coatings have been shown to exhibit low densities (
ρ ~ 3.3–4 g/cm
3, theoretical = 4.8–5.06 g/cm
3 [
29,
30]), which varies depending on coating microstructure [
31]. Buck [
31] observed the density of sputter-deposited amorphous coatings (
ρ ~ 3.3 g/cm
3) was lower than that of crystalline coatings (
ρ ~ 3.95 g/cm
3 for basally oriented coatings) due to the presence of microscopic vacancies and a high degree of disorder. Lince et al
. [
32] found that a higher oxygen content in MoS
2-xO
x coatings led to an increase in film density due to a reduction in crystallite size. Interactions between oxygen molecules and MoS
2 increased the defect density of the coating, thereby forming a disordered microstructure [
32]. Techniques such as ion-beam assisted deposition (IBAD) have been shown to improve both the density of pure MoS
2 coatings (
ρ ~ 4.4 g/cm
3) and wear resistance in humid and dry environments [
29].
A prevailing hypothesis for enhanced wear resistance of doped-MoS
2 is that it is linked to improved coating density [
23‐
28], though there is little to no direct evidence or a well-developed fundamental understanding of the role of this relationship present in the literature. A major barrier to this understanding lies in the difficulty in depositing fully dense pure films, or even films with consistent density. Variability in coating morphology is one of the main challenges limiting the widespread commercial use of sputtered pure MoS
2 films in engineering applications.
From a
research perspective, understanding temperature dependence, environmental resilience, or load-dependent friction behavior becomes even more challenging when coating microstructure varies from one deposition to the next. Relationships between detailed coating structure and deposition parameters such as substrate temperature [
33,
34], argon partial pressure [
31,
32], bias voltage [
37], and target-substrate distance [
8] have been studied for decades. While many parameters are controllable during deposition, there exist other variables and/or by-products that cannot. For instance, Buck [
38] studied the effects of water vapor in the plasma on the microstructure and tribological behavior of MoS
2 coatings. By varying the partial pressure of water, he showed that increased water vapor produced low density, less crystalline coatings that exhibited poor wear resistance. A secondary source of water vapor was also found to be a result of substrate heating during deposition via desorption of water from surfaces. The study found this source of contamination also led to less wear resistant coatings. Interestingly, substrate heating is often changed to improve the performance of MoS
2 coatings, yet it can have a negative effect depending on the cleanliness of the system.
From an applied perspective, sputtered MoS2 films are a preferred solid lubricant coating in inert environments, such as vacuum and space applications. However, despite the best efforts by coatings developers, batch-to-batch variations in film properties, which originate from difficulties in precisely controlling mutually influencing deposition parameters, pose significant challenges for hardware engineers owing to the resulting variability in functional behaviors, including tribological performance. Typically, aerospace hardware manufacturers provide witness coupons for each coating batch and subject them to in-house testing designed to qualify the batch. In recent years, several high-performance MoS2 based composites have fallen out of favor with hardware engineers due to the inability to reliably achieve the same caliber of tribological performance they once did. While the motivation for more wear resistant, lower friction, and environmentally agnostic materials always remains, there is an applied need to develop metrics to quantify what properties make universally “good” MoS2 coatings and encourage an understanding of what variables during or prior to deposition may be at play to disrupt this. This information will be invaluable in ensuring the quality and consistency of MoS2 films, as well as metrics that can enable future materials discovery and optimization of film composition and structure for a range of applications.
Given that even experienced commercial suppliers of sputtered MoS2 coatings can produce films with varying structure and performance due to uncontrolled deposition parameters, a method is needed to efficiently inspect coatings for critical attributes that will insure adequate tribological performance in the intended application. The purpose of this work is to show that density and hardness can be used as quality control metrics to insure the tribological performance of pure sputtered MoS2 coatings.
2 Experimental Methods
2.1 Materials Synthesis
Two manufacturers were asked to provide dense, nanocrystalline, pure MoS2 coatings in two separate deposition runs. These were requested to understand if (1) comparable coatings (i.e., orientation, porosity, tribological performance) could be made by separate manufacturers and, (2) if the same manufacturer could provide two identical batches of coatings. The samples include a “Low-Density Coating #1” designated here as LD-1, a “Low-Density Coating #2” designated LD-2, and a “High-Density Coating” designated HD-1; sample designations were based on the results of characterization presented later in this manuscript.
2.1.1 Deposition of LD-1
Pure MoS2 films were deposited in a vacuum deposition system (base pressure 5 × 10–6 torr) equipped with both radio frequency (RF) and direct current (DC) magnetrons. The substrates were affixed to a rotating stage that was biased at 50 VDC. The RF magnetron was used to sputter MoS2, and the DC magnetron was used to sputter a 99.99% pure titanium (Ti) target. The Ti target was sputtered first to create a ~ 100 nm thick Ti adhesion layer between the steel substrate followed by a ~ 200 nm gradient layer of Ti/MoS2 and then a ~ 800 nm thick pure MoS2 coating. Target powers were kept at 80 and 120 W, respectively.
2.1.2 Deposition of HD-1 and LD-2
Pure MoS2 coatings (~ 1 μm thick) were deposited on polished 440C steel substrates (~ 20 nm Ra roughness) with a 10 nm Cr adhesion layer (using arc evaporation) via DC magnetron sputtering using 1.5 mTorr Ar and a 3″ MoS2 target at 150 W and 30 V bias for 30 min. Identical processing conditions were used to produce HD-1 and LD-2.
2.2 Mechanical Test Methods
2.2.1 Hardness Measurements
Hardness values of MoS2 films were determined via nanoindentation on a Hysitron TI980 equipped with a Berkovich tip. Prior to experimental testing, the tip area function and load frame compliance were calibrated over the entire load range of the instrument with fused silica as the reference material. 5 × 5 indent arrays with a 10 µm spacing between indents were performed on each film. The maximum load in the load function was 1 mN. For each indent, a CMX (continuous measurement of X) load function was used, consisting of a constant strain-rate load superimposed with a 220 Hz oscillating load. The strain rate was kept constant at 0.123 s−1 to mitigate strain-rate effects and the oscillating load was employed to provide depth-dependent data. The instantaneous hardness H was calculated by H = Fmax/A, where Fmax is the maximum load and A is the contact area at each depth. H values were averaged over indentation depths between 40 and 100 nm to calculate the mean for each indent, as the mechanical properties in this regime were relatively constant. The reported value for each MoS2 film represents the mean and standard deviation from the 25 indents.
2.2.2 Tribological Testing
Tribological testing was performed on all three coatings simultaneously using a custom-built ball-on-flat high-throughput linear reciprocating tribometer in a controlled environmental chamber. A normal force of 1 N was applied to a 3 mm diameter 440C ball (~ 1 GPa Hertzian stress) on each sample by a load head consisting of a normal load cell connected perpendicular to the friction load cell (phidgets 100 g micro-load cell). The load head is then connected to a compliant titanium flexure driven by a stepper motor stage. The samples were mounted to a bidirectional linear reciprocating stage and tested at a sliding speed of 2 mm/s. Experiments were performed in a dry N2 environment (Mbraun Labstar pro, O2 < 0.5 ppm, H2O < 0.5 ppm) and air environments at 0%, 30% and 60% RH (± 2% RH) in a separate humidity-controlled enclosure.
2.3 Focused Ion Beam (FIB)/Transmission Electron Microscopy (TEM)
Cross-sections were prepared for transmission electron microscopy TEM analysis using a focused ion beam (FIB) in a Dualbeam ThermoFisher Helios. A 2 µm thick protective Pt layer was deposited by first the electron beam and then the ion beam to ensure the surface was not damaged by the FIB. The lamella was studied with the TEM at 200 kV (JEOL JEM-ARM200cF, Tokyo, Japan) and images were acquired with a Gatan Ultrascan CCD camera. Scanning TEM (STEM) dark-field and bright-field images were acquired with a probe size of 0.078 nm and the images were processed and analyzed in DigitalMicrograph (Gatan, Pleasaton, CA). TEM of LD-2 cross-sections was performed using an aberration corrected scanning TEM (FEI Titan™ G2 80–200 STEM) operated at 200 kV and high-angle annular dark-field (HAADF) imaging.
2.4 X-ray Diffraction (XRD)
A PANalytical Empyrean diffractometer with a Cu X-ray tube at a wavelength of 1.541 Å was used to take the X-ray diffraction (XRD) measurements. A Bragg–Brentano HD mirror with suitable slits were used to shape the incident beam to maximize irradiation on the sample. The diffracted beam was shaped with a 7.5 mm antiscatter slit and a soller slit and detected with a PIXcel3D-Medipix3 1 × 1 area detector in scanning line 1D mode. Symmetric θ-2θ (gonio) scans were taken with a step size of 0.0066° and counting rate of 25 s/step.
2.5 Rutherford Backscatter Spectroscopy (RBS)
Rutherford backscatter spectroscopy (RBS) was performed by Infinita Laboratories, Saratoga, CA. A beam of 1.9 MeV 4He + was used for RBS with detection at a 165° scattering angle. An average beam current of 5 nA with integrated charge of 2.0 µC was used. Coating densities were calculated from RBS areal densities (atm/cm2), RBS measured composition, and TEM measured thickness. The spot size of the RBS has a diameter of ~ 10 mm, or ~ 78.5 mm2 analysis region.
4 Discussion
Both the LD-2 and HD-1 coatings were manufactured in the same deposition chamber with nominally the same (controllable) deposition parameters, and by the same technician, albeit on different days. Additionally, the LD-1 coating was manufactured in a different chamber but with similar deposition parameters as LD-1 and HD-1. One of the most striking and quantifiable differences between the coating batches is the density. HD-1 (
ρ = 4.5 g/cm
3) has a density close to that of bulk MoS
2 (
ρ = 4.8–5.06 g/cm
3), exceeding the average density of IBAD coatings (
ρ = 4.4 g/cm
3) [
29]. Both LD-1 (
ρ = 3.55 g/cm
3) and LD-2 (
ρ = 3.04 g/cm
3) have measured density values that are well below HD-1, likely due to the formation of the voids observed in the TEM (Fig.
1).
Differences in density and void formation could be due to variations in crystallite orientation and degree of crystallinity, as indicated by XRD. Buck observed that crystalline pure MoS
2 coatings (
ρ ~ 3.8–3.95 g/cm
3) are denser than amorphous coatings (
ρ ~ 3.3 g/cm
3), and that increased basal-orientation improves density due to decreased porosity [
31]. Though our results show that LD-1, which is edge-oriented (as indicated by the (10
\(\overline{1 }\) 0) peak), is denser than the basally oriented LD-2 coating, both coatings have low peak intensities corresponding to their preferential orientations. Orientation and crystallinity can influence the friction and wear behavior of MoS
2, with highly crystalline, basally oriented coatings having lower initial friction coefficients and faster run-in to steady-state friction over amorphous microstructures [
39,
40]. Although nanocrystal-amorphous coatings have been reported to have lower wear rates than nanocrystalline coatings [
41], in this study it is challenging to decouple the individual effects of orientation and crystallinity from density on the tribological behavior. The impact of orientation on density and void formation is supported by the growth kinetics of MoS
2 during deposition. Low density, porous coatings are a result of the formation of edge-oriented MoS
2 crystallites providing reactive edge-sites for new deposits leading to a high vertical growth rate and decreased horizontal growth rate. As larger, vertically oriented lamellae form, they cause a shadowing effect, blocking incoming deposits resulting in the formation of voids [
42,
43]. Void formation can greatly impact hardness. This phenomenon is widely studied in other material systems such as ceramics [
44‐
49] where it has been observed that hardness increases as porosity decreases.
Pure MoS
2 coatings deposited by PVD are typically sub-stoichiometric with a deficiency in sulfur [
35,
50] and can have high levels of oxygen in the bulk (> 10%) [
32,
51]. Oxygen substituted into the crystal lattice of MoS
2 forms MoS
2-xO
x by substituting sulfur and results in a peak shift of the (10
\(\overline{1 }\) 0) due to a reduced lattice constant [
32]. Though the coefficients of friction observed for oxygen rich films are not as low as pure MoS
2 coatings, films containing high amounts of oxygen have been shown to have increased density from a reduced crystallite size, thereby producing lower wear rates than that of pure MoS
2 coatings [
32]. Addition of oxygen, which can be viewed as a dopant, means that “pure” is a misnomer for MoS
2 films without dopants because of the added benefits oxygen can impose. In this work, no statistically significant difference in oxygen content in the coatings under investigation was detected. Notably, the oxygen content was very low and close to the accuracy of the analytical tool. While we do not believe that oxygen is contributing to the improved densification of HD-1, it is not unlikely that sources of contamination during the deposition process could introduce unwanted oxygen or water. Sources such as adsorbed water on the deposition chamber walls due to exposure to lab air during sample changing or transfer and a contaminated sputtering target could be key uncontrolled factors that influence the coatings density and tribological behavior.
For MoS
2, relationships between coating porosity, hardness, and wear are not well established. Seynstahl et al
. [
52] varied the sample substrate rotation during deposition and observed that compact pure MoS
2 coatings with little to no porosity were harder (H = 5.69 GPa) and more wear resistant (k ~ 1 × 10
–7 mm
3/Nm) than softer (H ~ 0.06–0.25 GPa), porous films (k ~ 5 × 10
–6—2 × 10
−5 mm
3/Nm). Although the authors did not directly measure density, we observe a similar trend with HD-1 exhibiting a higher hardness (H = 4.4 GPa) and lower wear rate (k = 5.74 × 10
–8 mm
3/Nm) in dry N
2 than both LD-1 (H = 1.6 GPa, k = 7.98 × 10
–7 mm
3/Nm) and LD-2 (H = 2 GPa, k = 1.59 × 10
–6 mm
3/Nm). While density is a major factor contributing to the increased hardness of HD-1, the discrepancy between hardness and density for LD-1 and LD-2 could be due to differences in coating orientation. Though both LD-1 and LD-2 have weak peak intensities indicating nanocrystalline/amorphous microstructures, there is weak preferential vertical orientation of LD-1 (
i.e., basal planes parallel to the indentation axis), compared to the more basally oriented LD-2 (basal planes perpendicular to the indentation axis), allowing for deformation to occur between low shear strength basal planes as the coating is deformed. For vertically oriented coatings, the indenter tip can advance further into the coating by pushing the columns apart resulting in a lower measured hardness [
44]. The high hardness and density of HD-1 produced a greater wear resistance than that of widely established composite films such as MoS
2/Sb
2O
3/Au (k ~ 1 × 10
–7 mm
3/Nm [
21]) in dry N
2 environments. Improvements in the wear rates imparted by density and hardness are also observed in humid environments (Fig.
5), suggesting that wear performance measured in humid air could be a metric for a quality MoS
2 coatings, which for practical flight hardware, could be a useful metric if inert environments are unavailable or impractical to use.
Density as a driving factor for low wear MoS2 coatings, and hardness as an indicator of coating density, provides a useful metric for the qualification of MoS2 coatings to be used in practical applications. The low measured hardness of LD-1 and LD-2 would, for instance, indicate to an engineer that the batch of coatings will not meet specifications and should not be used. By using hardness as a metric, timely tribological testing of coating batches or costly characterization techniques such as RBS can be avoided, and only performed on batches such as HD-1 which meet an adequate hardness threshold. Additional useful metrics such as crystallinity and orientation, measured by XRD, would help provide a fast and accurate indication of a quality film.
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