Tribological Properties of 7A04 Aluminum Alloy Enhanced by Ceramic Coating

Abstract

The 7A04 Al alloy is a commonly used lightweight metal material; however, its low wear resistance limits its application. In this study, the wear resistance of this alloy was improved by preparing micro-arc oxidation (MAO) coatings, MAO/MoS₂ composite coatings, and hard-anodized (HA) coatings on its surface. The friction and wear behaviors of these three coatings with diamond-like coated (DLC) rings under oil lubrication conditions were investigated using a ring–block friction tester. The wear rates of the coatings on the block surfaces were determined using laser confocal microscopy, and the wear trajectories of the coatings were examined using scanning electron microscopy. The results indicated that, among the three coatings, the MAO/MoS₂ coating had the lowest coefficient of friction of 0.059, whereas the HA coating had the lowest wear rate of 1.47 × 10⁻⁶ mm/Nm. The MAO/MoS₂ coatings exhibited excellent antifriction properties compared to the other coatings, whereas the HA coatings exhibited excellent anti-wear properties. The porous structure of the MAO coatings stored lubricant and replenished the lubrication film under oil lubrication. Meanwhile, the introduced MoS₂ enhanced the densification of the coating and functioned as a solid lubricant. The HA coating exhibited good wear resistance owing to the dense structure of the amorphous-phase aluminum oxide. The mechanisms of abrasive and adhesive wear of the coatings under oil lubrication conditions and the optimization of the tribological properties by the solid–liquid synergistic lubrication effect were investigated. This study provides an effective method for the surface modification of Al alloys with potential applications in the aerospace and automotive industries.

1. Introduction

The 7A04 Al alloy is an important structural material in the aerospace [1] and automotive industries [2] owing to its low density, high specific strength [3], good ductility [4], and machinability [5]. However, it suffers from low hardness and poor wear resistance, which limit its application. To improve the wear resistance of Al alloys, a series of surface treatment techniques, such as ion implantation [6], laser treatment [7], and chemical conversion treatment [8], have been widely adopted in recent years. Although these technologies have made some progress, they fail to simultaneously provide the advantages of green environmental protection, low pretreatment requirements, simple processing, low environmental pollution, and excellent comprehensive performance of the film layer [9,10]. Therefore, in recent years, micro-arc oxidation (MAO), MAO/MoS₂, and hard-anodized (HA) coatings have attracted attention as advanced surface treatment technologies. These ceramic coatings are generated in situ on the surfaces of Al alloys to enhance the wear and corrosion resistance of the alloys, compensate for the insufficient film–base bonding force of the traditional technology, and optimize the wear resistance of Al alloys.

Hard anodizing is an electrochemical surface treatment technique that is usually performed in an acidic electrolyte such as sulfuric acid. A thick and uniform aluminum oxide film is generated on the surface of an Al alloy substrate by applying a high voltage and low current density [11]. Although this technique can significantly enhance the hardness of a material, the limited heat resistance of the HA film may lead to cracking or peeling in high-temperature environments, limiting its use in high-temperature applications [12]. In addition, the use of acidic electrolytes can cause environmental pollution.

MAO is a surface-strengthening technique for the in situ formation of a ceramic oxide film on the surfaces of workpieces made of metals such as Al, Ti, and Mg and their alloys [13]. By applying a voltage to the workpiece, an electrochemical reaction with the electrolyte solution is stimulated, resulting in a micro-arc discharge that produces a film with high densities and good mechanical properties. Conventional MAO coatings typically exhibit irregular undulations and cracks, resulting in uneven hardness of the coating surface, which affects its friction properties [14]. However, increasing the lubricity of MAO coatings, which can reduce friction, is possible. MoS₂, as one of the many-layered solid lubricant materials, is particularly suitable for severe working conditions owing to its high chemical stability, low coefficient of friction (CoF), and high load-bearing capacity [15,16]. The porous nature of the MAO coating provides dispersion space for MoS₂ to form an effective lubrication film during friction, reducing the CoF. The metal-based self-lubricating composite ceramic layer prepared according to this principle contains uniformly dispersed, phase-stable, and strongly bonded lubricating phases both internally and on the surface, and its friction-reducing effect has been confirmed in various application scenarios. Lv et al. [17] synthesized a MAO/MoS₂ composite lubrication layer via a hydrothermal reaction on the surface of the MAO coating and investigated the MoS₂ lubrication layer, which had high purity and a unique micro-spherical structure. Comparison of the friction performance of MAO/MoS₂ and MAO coatings under different loads involving ball-on-disk wear revealed that the lower CoF of the MAO/MoS₂ composite coatings was due to the self-lubricating properties of MoS₂ particles in conjunction with the formation of MoO₃ by the wear process. Chen et al. [18] prepared MAO composite coatings by adding graphene particles to the MAO electrolyte of the 6064 Al alloy and found that the addition of graphene particles changed the breakdown voltage, promoted the MAO reaction, and increased the thickness of the MAO coatings, which increased the coating hardness. Currently, most ceramic surface coating lubrication tests employ the ball-on-disk wear method, which, despite its popularity, is more susceptible to external factors such as temperature and load variations compared to line contact under point contact conditions [19,20,21,22,23]. Line contact is prevalent in mechanical systems, including sliding guides, cam followers, and cylinder and piston systems. Although ring–block wear tests have been utilized to simulate these real-world conditions, there is a lack of comprehensive understanding regarding the lubrication effectiveness of coatings and their mechanisms in line contact scenarios, which restricts the broader application of ceramic surface coatings in industry. In this context, MAO composite coating technology predominantly enhances the hardness and wear resistance of coatings by incorporating hard particles like SiC and Cr₂O₃ [24]. However, there is a scarcity of research focused on enhancing the wear resistance of MAO through the addition of lubricating elements. As the demand for electric vehicles grows, it becomes increasingly important to optimize the wear resistance and lubricity of MAO coatings to mitigate wear on critical components, prolong their service life, and decrease energy consumption, thereby fostering the sustainable development of the automotive industry.

In this study, the friction and wear behaviors of a friction pair consisting of the 7A04 Al alloy, a MAO/MoS₂ composite coating, a MAO-coated block specimen, and a diamond-like carbon (DLC) ring were investigated. Scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) were used to analyze the wear surface and elemental changes in the block specimens. Additionally, Raman spectroscopy was used to analyze the wear surface and friction wear mechanism.

2. Specimen Preparation and Testing

2.1. Coating Preparation

A block specimen of the 7A04 Al alloy, which nominal composition is presented in

Table 1. Chemical composition of the 7A04 Al alloy (wt.%).

Element	Zn	Mg	Cu	Cr	Si	Mn	Al
Content	5–7	1.8–2.8	1.4–2	0.1–0.25	0.4	0.2–0.6	Balance

, was employed in the experiment. After wire cutting a rectangular block of the size 19 mm × 12 mm × 12 mm, the specimen was sanded using 200#–1200# grit paper, rinsed with deionized water, ultrasonically cleaned in an acetone solution for 15 min at room temperature, and blown dry with cold air.

2.1.1. Micro-Arc Oxidation MAO

Prior to the MAO, the samples were treated with a NaOH and Na₃PO₄ pretreatment solution at 70–80 °C for 3–5 min. A basic alkaline electrolyte consisting of 10 g/L Na₂SiO₃, 5 g/L NaOH, and an appropriate amount of deionized water was used, and 4 g/L MoS₂ particles (grain size of 50 nm) were added to the basic electrolytes of the MAO/MoS₂ specimens. The added nanoparticles were prone to agglomeration; to prevent this, 4 g/L hydroxypropyl cellulose (HMC) was added to the base electrolyte and then ultrasonically stirred until it was uniformly suspended in the electrolyte. An asymmetric bipolar multifunctional alternating current pulse power supply was used, and the preparation parameters were set as follows: positive voltage, 460 V; positive frequency, 500 Hz; negative voltage, 60 V; duty cycle, 20%; oxidation time, 90 min; and current density, 3 A/dm². A constant current method was used to perform a MAO treatment on the surface of the Al alloy, and a cooling system was used to keep the temperature of the electrolyte below 45 °C. The electrolyte was then mixed via ultrasonic stirring until it was uniformly suspended in the electrolyte. During the MAO treatment, the sample was selected as the anode, and the electrolyzer was selected as the cathode. Following the treatment, the sample was rinsed with deionized water, blown dry with cold air, and placed in a sealed bag for subsequent testing and analysis. The detailed composition and parameters of the MAO electrolyte are presented in

Table 2. Electrolyte composition and process parameters for the MAO process.

Component	Value
Positive voltage (V)	460
Negative voltage (V)	60
Positive frequency (Hz)	500
Duty cycle (%)	20
pH	12.3
Oxidation time (min)	90
Current density (A/dm2)	3
Depositing temperature	≤45 °C
MAO Electrolyte	10 g/L Na2SiO3, 5 g/L NaOH, 4 g/L HMC
MAO/MoS2 Electrolyte	10 g/L Na2SiO3, 5 g/L NaOH, 4 g/L MoS2, 4 g/L HMC

.

2.1.2. Hard Anodizing—HA

HA coatings were performed in a properly prepared double-walled glass container. A direct current power supply with a constant current was used, with a current density of 2 A/dm². An aluminum block was used as the anode for the power supply, and stainless steel served as the cathode material. Ultrasonic stirring was employed to cool the electrolyte. The detailed composition and parameters of the HA electrolyte are presented in

Table 3. Electrolyte composition and process parameters for the HA process.

Component	Value
Voltage (V)	120
Frequency (Hz)	500
pH	1.2
Current density (A/dm2)	2
Oxidation time (min)	40
Depositing temperature	5 °C
HA Electrolyte	15% sulfuric acid

.

The ring specimens were made of 25Cr₃Mo₃NiNbZr mold steel and had an inner diameter of 42 mm, an outer diameter of 50 mm, and a width of 13 mm. The nominal composition of the material is presented in

Table 4. 25Cr₃Mo₃NiNbZr chemical composition (wt.%).

Element	C	Si	Mn	Cr	Mo	Ni	Nb	Zr	Fe
Content	0.28	<0.1	0.18	3.03	2.94	0.55	0.14	0.012	Balance

. The surface coating comprised DLC prepared using ionomer-enhanced chemical vapor deposition (PECVD), and the ring specimens were prepared using the same parameters. The chemical composition data in Table 1 and Table 2 were provided by the material manufacturers.

2.2. Microstructural Characterization

Field-emission SEM (JSM-IT800, JEOL Corporation, Tokyo, Japan) was used to examine the coating and cross-sectional morphologies of the specimen surfaces before and after wear. Elemental distribution analyses of the coating surface and wear marks were performed using EDS. The porosity of the coating surface under SEM was determined using ImageJ v1.54f software; three measurements were taken for each surface and averaged. The crystal structure was characterized using an X-ray diffractometer (Brux-D8, Bruker AXS, Karlsruhe, Germany) with Cu-Ka radiation at 40 kV and 40 mA. Cu was used as the target material, and the scanning range was 20°–80°, with a scanning speed of 6°/min. Hardness tests were performed using a microhardness tester (HVT-1000, Huayin Testing Instrument Factory, Laizhou City, China). This instrument was used to measure the microhardness of the surface at a certain depth with a load of 50 g and a loading time of 15 s. Owing to the surface roughness of the MAO film, the hardness test was performed after the surface of the sample was polished, the hardness of the structure of the uniform place of the 5-point hardness was tested, the mean value of the data points and the standard deviation of the data group was calculated, and a graph was made. A scratch tester (RST-300, Revetest, Peseux, Switzerland) was used to determine the coating and substrate bonding strength using Rockwell C-type diamond contacts (cone angle of 120° and contact radius of R = 200 μm). In the test, a diamond contact was applied to the coating surface under normal loads increasing from 0 to 100 N with a loading rate of 20 N/min and a scratch length of 5 mm. The critical loads were determined via specific damage observations and acoustic emission signal change analyses. The surface Raman spectra (InVia, Renishaw, London, UK) were examined using a confocal laser Raman spectrometer (LabRAM HR Evolution, laser wavelength of 532 nm).

2.3. Tribological Properties

A high-speed ring–block testing machine (MRH-3) was used under oil lubrication conditions at 25 °C with a normal load of 300 N, rotational speed of 200 rpm, rotation time of 20 min, and ring radius of 25 mm. The experiment was performed using the same substrate with three different coated block specimens. The ring was fabricated using the same process (DLC-coated 25Cr₃Mo₃NiNbZr mold steel). The blocks and rings were kept to maintain a smooth surface before the experiment, and lubricating oil was subsequently brushed onto their surfaces. A block-on-ring friction diagram is shown in Figure 1. At the end of the experiment, the blocks and rings were ultrasonically cleaned with anhydrous ethanol for 10 min and dried in an electric blast drying oven at 60 °C for 0.5 h. The worn-out specimens were then removed and weighed five times. Laser confocal microscopy (LCM) (0LS5100, Rui Ke System Integration Co., Ltd., Beijing, China) was used to measure the surface roughness of the coating and to model and analyze the wear surface topography. The wear rate was calculated via three-dimensional (3D) topographic modeling. Here, five cross-sections at equal distances from the abrasion mark trajectory were selected, the average cross-sectional area was calculated, and the average cross-sectional area multiplied by the length of the abrasion mark was used as the abrasion mark volume V. The abrasion rate was calculated using the following equation:

$$\mathrm{Wear}\mathrm{rate}={\displaystyle \frac{V}{Fn\times S}}\left({\mathrm{mm}}^3/\mathrm{Nm}\right),$$

(1)

where V represents the wear volume (mm³), Fn represents the normal load (N), and S represents the total friction distance (m).

3. Results and Discussion

3.1. Analysis of Microscopic Morphology Before Wear

Figure 2 presents the micromorphology of the surfaces of the MAO, MAO/MoS₂, and HA coatings on the 7A04 Al alloy before wear (backscattered electron (BSE) diagram). As shown in Figure 2a, the MAO coating surface mainly consisted of cracks, micropores, and raised cellular structures. Figure 2b indicated that, for the MAO/MoS₂ coating surface, cellular projections were the main features. The surface contained a few cracks, and the irregular cellular projections contained micropores. The formation of these micropores is related to the generation of gases from the molten oxide eruption and the generation of discharge channels. Comparing the microscopic morphologies of the MAO and MAO/MoS₂ coatings revealed that fewer cracks and micropores existed in the MAO/MoS₂ coatings. Additionally, the surface porosity of the MAO coatings was 10.221%, whereas that of the MAO/MoS₂ coating was 5.638%, as measured using ImageJ. The main reason for the large difference in surface porosity between the coatings is that the addition of MoS₂ nanoparticles increased the conductivity of the electrolyte, resulting in a more molten matrix ejected from the pores and oxidation. The micropores and cracks generated during the growth process were adsorbed on the coating pores and sealed with the diffusion and electrophoretic deposition of MoS₂ nanoparticles, which increased the surface density [25]. The surface of the HA coating (Figure 2c) exhibited small pits and a few holes and cracks on the surface, and the surface oxide film was less porous. The HA coating had the lowest porosity of 2.773%, as measured using ImageJ. Figure 2d shows the coating base material (7A04 Al alloy). A few cracks were observed on the surface. These cracks were due to the low plasticity of the 7A04 Al alloy, which led to the accumulation of internal residual stresses during processing.

Table 5. SEM images of the coated surfaces corresponding to elemental scans of (a) MAO coating (b) MAO/MoS₂ coating, and (c) HA coating.

Element	Content, wt.%
	a	b	c
O	44.15	44.87	54.21
Al	32.08	32.16	40.21
P	2.03	2.39	-
Si	15.96	17.85	-
S	-	0.43	5.13
Mo	-	0.22	-
Others	4.62	2.08	0.53

presents the results of elemental analysis using the micromorphology energy spectra of the coatings shown in Figure 1. The ceramic coating was mainly composed of Al and O, the MAO coating surface contained Si and P, and a small amount of MoS₂ in the MAO/MoS₂ coating was deposited on the coating surface. The MAO coating surface has a high Si content, because the amorphous SiO₂ formed in the electrolyte cannot be dissolved in the electrolyte; thus, it was deposited on the coating surface as the coating grew [26].

Figure 3 presents cross-sectional views of the MAO, MAO/MoS₂, and HA coatings and the corresponding elemental line sweeps. The measured average thicknesses of the MAO, MAO/MoS₂, and HA coatings were 60 ± 2.52 μm, 60 ± 3.79 μm, and 30 ± 1.12 μm, respectively. After polishing of the cross-section, as shown in Figure 3a,b for MAO and MAO/MoS₂, respectively, the undulation of the coating interfaces and surface contour lines were clearly observed, indicating that the coating grew inward and outward simultaneously and that the inward growth could increase the bonding force between the coating and the substrate [27]. In Figure 3b, the MAO/MoS₂ cross-section is denser, thicker, and has fewer cracks than MAO. The reason for the denser and thicker coating was related to the addition of MoS₂ to increase the discharge voltage and deposition [28]. The cross-sectional line sweeps of MAO and MAO/MoS₂ presented in Figure 3(a₁) and Figure 3(b₁), respectively, show a consistent rise and fall in the O and Al contents, respectively. The Si content gradually decreased from the outside to the inside, whereas the P content exhibited the opposite distribution, a phenomenon that was closely related to the pretreatment of the samples. Pretreatment with the phosphate electrolyte promoted the rapid formation of a dense ceramic film on the substrate surface. This step had an important influence on the initial formation stage of the coating, enhanced the bonding of the film layer with the substrate, and improved the densification of the coating. As shown in Figure 3c, the HA cross-section had a small number of holes and cracks, with no obvious cracks at the interface, and the contour line of the coating surface had a small undulation. As indicated by the three cross-sections, these coatings bonded well with the substrate.

Table 6. Characteristic properties of the three coatings.

Sample	Thickness/μm	Porosity/%
MAO	60 ± 2.52	10.221
MAO/MoS2	60 ± 3.79	5.638
HA	30 ± 1.12	2.773

summarizes the characteristics of the three coatings.

The phase structure of the elements in the coating was analyzed, and the results are presented in Figure 4. As shown, the three coatings can be detected in the high-intensity Al peaks, indicating that the high-intensity X-rays passed through the thick coating into the substrate. The diffraction peak shapes of the HA coating and the 7A04 Al alloy can only be observed in the Al peaks, indicating that the hard anodization of the oxide film layer after the generation of Al₂O₃ resulted in an amorphous phase. For the MAO and MAO/MoS₂ coatings, the common phases of α-Al₂O₃ and γ-Al₂O₃ were detected, of which the main phase was γ-Al₂O₃. The X-ray diffraction (XRD) pattern of the MAO/MoS₂ coatings exhibited an additional α-Al₂O₃ peak and a mullite peak. This formation was due to the gradual isolation of the initially formed dense layer from the electrolyte as the oxidation time increased, which reduced the cooling rate. The internal γ-Al₂O₃ was gradually converted into α-Al₂O₃ under the effect of arc heat [29]. Meanwhile, the growth of MAO coatings in the silicate electrolyte was mainly manifested by the deposition of SiO₂, which was accompanied by the mild oxidation of the Al substrate. With an increase in the voltage, the generated Al₂O₃ and SiO₂ layers gradually melted and sintered, forming a mullite structure. For the samples coated with MAO/MoS₂, the MoS₂ peaks were located at 40° and 59°, confirming that the MoS₂ particles participated in the MAO reaction and were successfully doped into the coating.

3.2. Mechanical Properties of the Coatings

Figure 5 shows the microhardness of the MAO, MAO/MoS₂, and HA coatings and the substrate. The hardness of the Al alloy was significantly increased after the formation of coatings on the surface. The hardness of the MAO coating was 483.7 HV, which was lower than that of the added MAO/MoS₂ coating. The average hardness of the MAO/MoS₂ coating was three times that of the 7A04 Al alloy. The HA coating had the highest average hardness, which was 4.1 times that of the substrate. The average hardness of the 7A04 Al alloy was the lowest (167.8 HV). The hardness of the coatings was affected by the phase composition and microstructure. The XRD results indicated that the MAO/MoS₂ coatings had a mullite phase and more α-Al₂O₃ phase, and the cross-sectional plots indicated that the MAO/MoS₂ coatings had a larger average thickness, which increased the average hardness of the coatings [30,31].

The RST-300 instrument was used to determine the bond strength between the ceramic coatings and substrate. Figure 6a–d present the acoustic emission signal versus friction curves of the MAO, MAO/MoS₂, and HA coatings and the corresponding scratch morphologies. Owing to the ceramic particles in the MAO and HA coating layers and the influence of the surface quality, the testing process generated interfering signals that affected the accuracy of the critical loads [19,32]. Therefore, acoustic signals and friction curves were used in conjunction with scratch morphology maps (BSE images) to determine the critical load of the bonds (Lc₂). In the plots of the acoustic emission signals and friction profiles in Figure 6a–c, the friction profile of the MAO coating varied slowly, while the friction profile of the HA coating varied more rapidly. In Figure 6a,b, the MAO coating exhibited a few small peaks up to 70 N and started to exhibit significant peaks only above 70 N. Figure 6c shows a distinct peak corresponding to the HA coating acoustic signal at 25 N. To determine the exact Lc₂ value, the scratch morphology was examined using SEM. Figure 6d shows that the scratches are mainly divided into two parts: the grey ceramic layer in the first half of the abrasion mark and the silver substrate in the second half. In Figure 6d, the HA coating exhibited scratches in the grey–silver junction (2.3 mm; that is, a pro-bonding force of approximately 43 N). The acoustic signal at approximately 30 N appeared as an interfering signal that may be caused by the stripping of the membrane layer. The results indicate that the synthesis can be performed when the HA coating Lc₂ is 43 N. The acoustic signal bursts of the MAO and MAO/MoS₂ coating film layer failure matched the interface junction of the scratch micrograph in Figure 6d; thus, it was determined that the Lc₂ values of the MAO and MAO/MoS₂ coatings were 73 and 78 N, respectively. Observation of the edges of the scratched tracks did not reveal large-scale flaking. The resulting debris was scattered within the abrasion marks, and the coatings all exhibited good bonding strength.

3.3. Tribological Properties of the Coatings

The CoF versus sliding time curves for the MAO, MAO/MoS₂, and HA coatings and the 7A04 substrate under an applied load of 300 N are presented in Figure 7. As shown in Figure 7a, the initial CoF of the MAO coating was approximately 0.14. The CoF then decreased rapidly to 0.9, increased again, and gradually stabilized between 0.07 and 0.08 in 400 s. The fluctuation in the CoF between 50 and 200 s may have been due to the presence of large abrasive particles between the contact surfaces during the wear process. As the wear process proceeded, these abrasive particles were continuously extruded under a high contact pressure to form small particles that could eventually be stably embedded and stored between the contact surfaces. This dynamic change led to a corresponding fluctuation in the CoF. In Figure 7b, the MAO/MoS₂ and MAO coatings exhibited similar CoF change trends. The CoF at the beginning of the wear was approximately 0.13. After a relatively short period of time, to enter the stable wear stage, from the friction coefficient graphs, it could be observed in the wear periods of 20 and 100 s around the CoF fluctuation, which might be related to the abrasive particles, and then the stable wear stage was entered, and the CoF stabilized at approximately 0.5–0.6. Compared to the MAO coating, the addition of MoS₂ significantly reduced the wear time and resulted in a lower CoF in the stabilization stage, suggesting that the MoS₂ deposited on the MAO coating played a self-lubricating role. As indicated by Figure 7c, the CoF of the HA coating was 0.23 at the early stage of wear, and it stabilized at 0.12 after a short period of time and returned to 0.12 after a short fluctuation of approximately 700 s until the end of wear. Owing to the localized areas of the HA coating with bumps (e.g., Figure 7c), the oil film failed to cover the contact surface at the early stage of wear, resulting in a higher CoF at this stage. After the surface bumps were worn out, the surface morphology gradually became smooth, and the CoF stabilized. Figure 7d shows that the initial CoF of the Al alloy substrate was 0.15 and that the CoF decreased continuously until the end of wear.

The wear rates and average CoFs of the MAO, MAO/MoS₂, and HA coatings and 7A04 substrate after wear are presented in Figure 8. In Figure 8a, the 7A04 substrate exhibited the highest wear rate, indicating that the coating had a protective effect on the substrate. The wear rate of MAO/MoS₂ was 13.73 × 10⁻⁶ mm³/Nm lower than that of MAO (17.28 × 10⁻⁶ mm³/Nm) owing to the self-lubricating effect of MoS₂. The HA coating exhibited the lowest wear rate. Notably, it also had the highest hardness and highest CoF owing to the generation of amorphous-phase aluminum oxide, which was relatively dense and had few cracks, thus exhibiting better wear resistance. The above experimental phenomena indicated that all the stabilized CoFs of the 7A04 Al alloy treated with MAO were low. Figure 8b shows that the average CoFs of the MAO, MAO/MoS₂, and HA coatings and the 7A04 substrate were 0.077, 0.059, 0.124, and 0.069, respectively. Owing to the porous nature of the MAO coating, the lubricant could be stored so that a stable oil film could be formed on the contact surface to protect the surface. The lowest CoF for the MAO/MoS₂ coating suggested that the MoS₂ slowed the increase in the roughness of the coating surface and formed a more stable oil film on the contact surface, resulting in a lower CoF, which reduced the wear rate [33,34].

Figure 9a–d show the 3D morphology of the abrasion marks of the MAO, MAO/MoS₂, and HA coatings and the 7A04 substrate and the corresponding cross-sectional curves (Figure 9(a₁–d₁)). To examine the CoF changes during the break-in period, the surface and substrate roughness of the three coatings were measured using LCM, yielding the following results: MAO, 4.67 μm; MAO/MoS₂, 7.27 μm; HA, 2.06 μm; and 7A04 Al alloy, 0.977 μm. As shown, the widths of the abrasion marks of the MAO and MAO/MoS₂ coatings were divided into 4845.8 and 4252.5 μm, the HA coating had the narrowest abrasion mark (2004.4 μm), and the 7A04 substrate had the widest abrasion mark (4845.8 μm). In the cross-sectional diagrams (Figure 9(a₁–d₁)), the abrasion depth of the MAO coating was 95 μm, the abrasion depth of the MAO/MoS₂ coating was smaller (88 μm), and the HA coating had the smallest abrasion depth (18 μm). The maximum depth of the abrasion marks on the 7A04 substrate was 168 μm. The results indicated that all the coatings played a protective role with regards to the abrasion of the 7A04 substrate, and the wear resistance of the 7A04 Al alloy was increased. The HA coating exhibited the highest abrasion resistance, which was correlated with its surface roughness and densification.

3.4. Wear Surface Morphology Analysis

To investigate the wear mechanism of the coating, SEM characterization and the corresponding elemental analysis of the micromorphology of the abraded surface were performed, as shown in Figure 10. Figure 10a presents the grooves and black flaky bonded lumps on the wear scar surface. In Figure 10b, there were inconspicuous grooves and numerous black flaky bonding blocks on the surface of the abrasion marks, with localized coating detachment and obvious grooves on the substrate surface. Figure 10c shows the abrasion scar surface with evident grooves and a few black bonding blocks. Figure 10d shows the wear scar surface with obvious grooves, flaking, and cracks. In the EDS diagrams of the wear, the black bonding masses on the wear scar surfaces corresponded to the element C, which came from the DLC coating on the dyad. From the O elemental spectra, obvious parallel grooves were observed between the HA coating and the substrate. The grooves on all the wear marks were parallel to the sliding direction, indicating that no uneven load distribution existed during the wear process. According to the aforementioned characterization results for the wear marks, the main wear mechanism of all the specimens was a combination of abrasive and adhesive wears.

The wear morphology was further analyzed, and a comparison revealed that fewer grooves existed on the surface of the MAO coating, indicating that, in the early stage of ring wear, the MAO coating had a high roughness, direct contact between the friction partners occurred, and the actual contact involved a localized protrusion of the coating in contact with the opposing ring. Under a high load and cyclic shear stress, the hard debris wrapped around the coating was adsorbed on the surface of the coating with high hardness, which constituted three-body abrasive wear. When hard debris moved on the surface or inside the transfer film, it caused plastic deformation of the coating surface as it rolled. When the wear was further aggravated, the hard abrasive particles slid and cut the coating, resulting in deep furrows on the surface of the abrasion marks of the coating. As shown in Figure 10b, the surface of the abrasion marks had more transfer film and fewer abrasion marks, and the surface roughness of the MAO coating increased with the addition of MoS₂. Owing to its uneven surface properties, the contact area between the localized protrusions of the coating and the dyad was reduced during the initial stage of friction, significantly increasing the pressure at the contact point. This uneven microscopic contact surface led to the conversion of mechanical energy into thermal energy at the contact point, which increased the temperature of the contact area. The increased pressure and temperature of the contact surface may intensify the graphitization process on the DLC coating surface, increasing the tendency for adhesive wear to occur during the wear process. Jin et al. [35] also reported that severe adhesive wear occurred on the grinding balls with MAO/DLC coatings. Its surface energy spectrum indicated the presence of S, suggesting that MoS₂ played a lubricating role in the friction process and that forming a protective lubricating film with the dyadic parts under the action of shear force was easy, which played a lubricating role. This was reflected in the minimum CoF throughout the friction process (e.g., Figure 7b), and the surface of the wear marks had fewer grooves. Additionally, localized coating flaking was observed (Figure 10b). This may be due to microdefects in the MAO coating, which were prone to form crack sources under high loads. During the sliding test, these crack sources expanded, and spalling occurred under periodically applied high loads. In Figure 10c,d, the abrasion patterns were mainly plow grooves parallel to the sliding direction and a small amount of transfer film. The debris generated during the abrasion process acted on the contact surface to form the main abrasion patterns. A small amount of C was detected via EDS, as shown in Figure 10c, indicating that the stable wear phase accounted for a major part of the wear process.

Raman spectroscopy was performed on the wear marks to further analyze the influence of the friction process. Figure 11 shows the Raman spectra of the MAO-, MAO/MoS₂-, and HA-coated wear mark areas. A Gaussian function was fitted to the range of 400–2000 cm⁻¹ to obtain the D peak associated with the graphite phase near 1360 cm⁻¹ and the G peak associated with the sp² C–C bond near 1580 cm⁻¹. The variation in the I_D/I_G value can be used to analyze the effect of the transformation of DLC into a graphite structure during sliding in the friction process [36]. As shown in Figure 11d, the maximum I_D/I_G value was 1.08 for the MAO/MoS₂ coating. This indicated an increase in the transformation of the DLC ring surface in contact with the MAO/MoS₂ coating into the graphite structure. With an increase in the degree of graphitization, more C-based transfer films were generated, and the generated C-based transfer film entered the oil film and was stored together with the MoS₂ particles, which provided anti-friction and anti-wear effects.

MAO coatings can form thick oil films during wear owing to their inherent porosity. These oil films not only enhance the lubrication effect but also significantly improve the ability to capture exfoliated C-based transfer films [35]. In particular, larger flake transfer films are more easily captured and carried by oil films because of their larger surface areas and lower surface energies. As the wear process continues, the oil film decreases, making it difficult for large flake films to exit the system with the lubricant. Consequently, more C-based transfer films are observed to adhere to the wear surfaces of the MAO coatings, particularly the MAO/MoS₂ composite coating [37].

4. Conclusions

In this study, MAO, MAO/MoS₂, and HA coatings were prepared on the surface of the 7A04 Al alloy. The microstructures of the three coatings, as well as their tribological properties and wear mechanisms under oil lubrication conditions with DLC ring–block wear, were investigated. The following conclusions were drawn.

(1) The surface of the 7A04 Al alloy was effectively modified by the MAO, MAO/MoS₂, and HA treatments, which significantly improved its hardness and wear resistance.

(2) Among the coatings tested, the MAO/MoS₂ coating exhibited the lowest CoF, with superior anti-friction performance, owing to the self-lubricating properties of MoS₂. Meanwhile, the HA coating exhibited superior wear resistance owing to the dense structure of the amorphous-phase aluminum oxide.

(3) Under oil lubrication conditions, the main wear mechanisms of the MAO and MAO/MoS₂ coatings were abrasive and adhesive wear, whereas the wear mechanisms of the HA coatings and 7A04 Al alloy were dominated by abrasive wear, accompanied by slight adhesive wear. The surface micropores of MAO coatings store oil, which can replenish the lubricating film of the liquid and impede the movement of wear particles. In addition, the MoS₂ particles and the solid lubrication provided by the graphitized layer work together to achieve a synergistic solid–liquid lubrication effect with the liquid lubrication film, optimizing the tribological properties.