Plasma electrolytic oxidation (PEO) [[1], [2], [3], [4], [5], [6]], as an eco-friendly surface modification technology based on plasma-assisted anodic oxidation for lightweight materials (aluminum, magnesium, titanium etc.), allows the creation of ceramic-like oxide coatings with firm adhesion to the substrate. Previous studies have demonstrated that PEO coatings, particularly alumina coatings, possess excellent resistances to environmental corrosion and abrasion damages, thus, this technique has been applied on aluminum alloys in a wide range of industrial applications for decades [[7], [8], [9], [10], [11]]. Many studies [[12], [13], [14]] focused on optimizing the combination of electrolytes and electric parameters in order to improve the performance of resultant PEO coatings. However, the quality of PEO coatings is often determined also by the chemical composition of the aluminum alloys, either by the alloying elements influencing PEO layer composition or via formation of secondary phases affecting the PEO process directly [15,16].

To investigate the effects of chemical composition of substrates, attempts have been made to study aluminum alloys, which contain different alloying elements. The major alloying element Cu generally exists in form of precipitates in aluminum alloys. An earlier study [17] explained the absence of Cu in PEO coatings of a commercial Al-Cu-Mg alloy (3.8-4.9 wt% Cu) by the dissolution of Cu-rich precipitates into the electrolyte during PEO processing and incapability of being absorbed and deposited into PEO coatings. However, a later research [18] reported the irregularity distribution of Cu in the outer layer of PEO coating on 2024 aluminum alloy, confirming Cu incorporation in the coatings. Afterwards, Oh et al. [19] identified approximately a Cu residual ratio of 0.1 in the PEO coating relative to that in the substrate of 2024 aluminum alloy and revealed an adverse role of Cu in α-Al₂O₃ transition. A brownish-red PEO coating on 2A97 aluminum alloy (3.8 wt% Cu) was obtained, which was attributed to the generation of dark red Cu₂O from oxidation of Cu in 2A97 aluminum alloy [20] by specific B-type discharges [21]. Furthermore, Veys-Renaux et al. [22] tracked the evolution of Cu-rich precipitates in AA2214 aluminum alloy during PEO processing and discovered the de-alloying of Cu-rich precipitates into Cu nanoparticles. Cu nanoparticles increased the electrical conductivity of the oxides, delaying the initiation of dielectric breakdown.

PEO coatings of binary Al-Mn alloys with increasing alloying content of Mn (1-8 at.%) exhibited increased surface roughness, coating thickness, amount and size of porosities [23]. Electrochemically active Mn, having a similar oxidation potential to Al, could easily substitute Al in Al₂O₃ crystal structure thus accelerating coating growth. The higher vapor pressure of Mn resulted in a higher porosity and hence higher surface roughness of PEO coatings during spark formation. In PEO studies of 2214-T6 and 7050-T74 aluminum alloys, the inhibition of α-Al₂O₃ growth in the coating of 7050-T74 aluminum alloy was ascribed to the abundant Zn in the substrate [24]. However, another study [19] identified a more significant contribution of Mg to suppress the formation of α-Al₂O₃ as compared with that of Cu and Zn. PEO coating formed on high Si containing aluminum alloys were mainly composed of ɣ-Al₂O₃, α-Al₂O₃, SiO₂ and a small amount of mullite phase (Al₂O₃·SiO₂) [25]. The presence of Si in Al matrix led to lower efficiency in coating formation and lower porosity of the coating during PEO treatment due to the lower electrical conductivity of Si [26].

Although the influence of separate alloying elements in the substrate on the characteristics of resultant PEO coatings has been investigated, there is still a challenge to control the synthesis of PEO coating when there is lack of knowledge about coating formation and evolution on different materials and phases during PEO processing for multiphase aluminum alloys. AlSi9Cu3 alloy as an engineering material normally requires surface pretreatment before application [[27], [28], [29]]. Unfortunately, the large-sized secondary phases and high quantity of Si restrict the conventional anodization since the thinner and vulnerable anodic film is not able to fulfill the mechanical and tribological requirements [22,26]. As a potential solution, PEO treatment was performed on AlSi9Cu3 alloy in this work. It is noteworthy that AlSi9Cu3 alloy is a good model system of hybrid materials for the purpose of investigating PEO coating growth on different materials (α-Al, Si) and phases (Al₂Cu and β-Al₅FeSi) at the same time. Our studies disclosed the inhomogeneous growth and discontinuity of PEO coatings in presence of these different phases in AlSi9Cu3 alloy. The growth behavior of the coating from initial stage of PEO coating formation and further coating evolution was analyzed in detail and discussed.

Rectangular specimens of as-cast AlSi9Cu3 alloy with dimensions of 15 mm × 15 mm × 4 mm were selected. The mass fraction of the substrate is 9.57 wt% Si, 2.56 wt% Cu, 0.93 wt% Sn, 0.83 wt% Mg, 0.38 wt% Zn, 0.35 wt% Fe, 0.14 wt% Mn and Al balance. Prior to PEO treatment, the surface of the specimens was ground using abrasive papers from 1200 grit to 4000 grit and polished with 3 μm polishing agent, then flushed by ethanol and distilled water respectively and dried in air at ambient temperature.

PEO process was carried out by using a unipolar-pulsed DC power source. In this work, a constant current regime was applied with a current density of 3 A/dm² and a duty cycle of 10% (t_on : t_off =0.5 ms : 4.5 ms) in an alkaline electrolyte (pH = 12.4, conductivity 23.2 mS/cm) composed of 20 g/L Na₃PO₄ and 1 g/L KOH. The test specimens connected with a conductive holder worked as anode. The stainless steel tube being part of the heat-exchange system was adopted as cathode. The temperature of electrolyte was kept at 20 ± 1 °C. The PEO treatment was implemented at different treatment times ranging from 15 s to 480 s (15 s, 30 s, 60 s, 120 s, 240 s and 480 s, respectively).

The resulting voltage as a function of processing time was recorded with a data acquisition system (SignaSoft 6000 software package) provided by Gantner. The distribution of different phases of the substrate, surface and cross-sectional morphologies as well as chemical composition of the coatings were evaluated by means of scanning electron microscope (SEM, TESCAN Vega3 SB) equipped with an energy dispersive spectrometer system of eumeX (EDS, IXRFsystems). The phase composition of the substrate and coatings were analyzed using X-ray diffraction (XRD, D8 Advance, Bruker AXS) with CuK_α radiation (40 kV, 40 mA). Diffraction pattern were obtained at diffraction angle 2θ from 20° to 80° with a step size of 0.02°/s and grazing angle of 1°.

Fig. 1 reveals that AlSi9Cu3 alloy is mainly composed of α-Al and Si with additional intermetallic compounds Al₂Cu and β-Al₅FeSi respectively. The characteristic peaks of intermetallics can be hardly observed because of their lower intensity relative to the matrix. Simultaneously, Fig. 2 illustrates the microstructure of the bare AlSi9Cu3 alloy with the phases indexed, in accordance with the study of Puga et al. [30]. The microstructure is characterized as grey α-Al matrix with lath-liked eutectic Si, herringbone-structured intermetallic Al₂Cu and irregular intermetallic β-Al₅FeSi. At higher magnification (Fig. 2(c) and (d)), the overlapping distribution of intermetallic Al₂Cu and β-Al₅FeSi is revealed as another feature of the substrate microstructure. The observation of substrate microstructure provided visual information to differentiate the phases and further assisted in monitoring the coating evolution on these phases during PEO processing.

3.2.1. Voltage evolution during PEO processing

Fig. 3 shows the change of voltage with increasing treatment time until 480 s. It is the typical form of voltage-time curve when a constant current density is applied in PEO processing of aluminum alloys [31,32]. The curve can be divided into two discernable stages: firstly, the stage of conventional anodic oxidation featured by a rapid and linear increase of oxidation voltage at a rate of 3.3 V/s, when the treatment time is shorter than 60 s. The first stage indicates the initial mostly electrochemical-driven dissolution of the substrate and the formation of thin passive barrier layer on the interface of substrate/electrolyte [33]. First tiny arcs were observed after 60 s at a voltage of 200 V. Afterwards, the process progresses into the second stage, characterized by the phenomenon of numerous visible micro arcs moving around on the surface of the specimen. However, at this point, the slope of the voltage-time curve decreases associated with the continuous thickening of the porous PEO coating. The PEO process remains at this stage until 480 s and the increase of voltage is slowing down further, reaching a final voltage of 350 V.

3.2.2. Phase composition

The phase composition of the coatings after different treatment times is shown in the XRD patterns (Fig. 1). Besides the intensive peaks from the substrate (α-Al, Si and Al₂Cu) due to the X-ray penetrating through the coating into the substrate, some weak signals (mainly after 480 s treatment) demonstrate that the coatings consist of γ-Al₂O₃ and mullite (3Al₂O₃·2SiO₂) respectively. Interestingly, although the phase composition of the coatings seems not to change with treatment time, some substrate signals appear to be stronger indicating that in spite of coating formation the Al matrix is preferentially etched with intermetallic Al₂Cu remaining. Only at 480 s, an additional amorphous phase appears as a broad bump at 2θ around 20° to 35°. The transformation from amorphous alumina to crystalline Al₂O₃ mostly depends on the temperature [34]. However, discharges providing enough energy for phase transformation and heating the specimen and electrolyte up, are still not occurring within the short treatment times and a final voltage of 350 V.

3.2.3. Surface morphologies and element distribution

Fig. 4 displays the surface morphologies on AlSi9Cu3 after different PEO treatment times and reveals the coating evolution on the four different phases (α-Al, eutectic Si, Al₂Cu and β-Al₅FeSi intermetallics). At the very beginning (the first 15 s), initial selective dissolution occurs in the interface to Al₂Cu, as the denuded net structure of Al₂Cu is clearly revealing in Fig. 4(a). In the meantime, the deposition of conversion products is observed preferentially surrounding the periphery of the intermetallic Al₂Cu, which is supported by O and P signals shown in the EDS analysis in Fig. 5(a).

After 30 s, the conversion products accumulate around Al₂Cu and β-Al₅FeSi intermetallics as revealed in (Fig. 4(b)). The signals of the characteristic elements Fe, O and P in these positions (Fig. 5(b)) confirm the involvement of β-Al₅FeSi in formation of the conversion products. However, the coverage of the surface in larger distance from these intermetallics by elements O and P originated from the electrolyte is limited and the boundaries of Si are still recognizable. The conversion products pile up further until 60 s treatment (Fig. 4(c)) and the elements O and P start to cover the entire surface (Fig. 5(c)), whereas, it seems that the conversion products peel off from the surface, leaving behind some voids around the Al₂Cu phases after 120 s (Fig. 4(d)). According to observation of the sparking, the first sparks appear already after 60 s and it is likely that they are involved in the removal of the conversion products around the intermetallics.

The topography of intermetallic Al₂Cu after 120 s features some lower areas (Fig. 6(a)), which are obviously a result of the preferential dissolution of Al₂Cu intermetallics. Micro-pores from initial discharges can be seen at 120 s as well. They are largely distributed on the junction of intermetallic Al₂Cu and surrounded α-Al matrix (Fig. 6(a)) and less on the surface of Al₂Cu and β-Al₅FeSi (Fig. 6(b)). It can be inferred that the initial discharges have started on certain regions of the barrier layer when the voltage reached 200 V after 60 s. These regions are apparently related to the fast growing conversion layer on the top of α-Al matrix around the intermetallics. When the voltage gradually ramps up to 250 V at 120 s, it also exceeds the breakdown potential of the barrier layer on all phases except of eutectic Si [35], because of its lowest rates in anodizing and formation of oxides illustrated by the weak coverage of O/P on it (Fig. 6(c)).

The surface is firstly fully covered with a typical PEO coating showing numerous micro-pores after 240 s treatment (Fig. 4(e)). The continuous coverage of O on eutectic Si (Fig. 7(a)) reveals that silicon surface is oxidized and an oxide layer is formed already. Interestingly, there are discernible surface micro-morphologies distinguishable by the distribution of micro-pores of different sizes. According to the distribution of the characteristic elements (Al, Si, Cu and Fe), the three distinct surface micro-morphologies are related to the phases present at the surface, indicating the substantial effect of secondary phases on the surface morphologies of the coating at this stage. At higher magnification, the coating on eutectic Si (Fig. 7(b)) shows a few small micro-pores and the boundaries of eutectic Si underneath are still visible because the layer is too thin to cover the morphology of eutectic Si completely. The mostly sintered coating with the largest micro-pores is certainly growing on the overlapping area of Al₂Cu and β-Al₅FeSi in Fig. 7(d).

After longer treatment (480 s), the surface morphologies of PEO coatings evolve significantly (Fig. 4(f)). Besides micro-pores spread on the entire surface, even some micro-cracks form already in particular regions at 480 s. An increase in micro-pore sizes is observed on the PEO coating surface compared with that of 240 s (Figs. 7(a) and 8 (a)) due to the higher energy (voltage and current) in the discharge as resistance of the coating is increasing with thickness of the coating. Three micro-morphologies characterized by divergent porosities as observed on the coating surface of 240 s are still discernible with extended treatment time, shown in Fig. 8(b)-(d)) (marked areas 1, 2 and 3). The inhomogeneity in the micro-morphology is mainly controlled by the intensity and density of discharges, which can be related to different resistance of the oxide film on the different phases. The main element composition and content of the coating in the three different micro-morphologies are determined in Fig. 8(e). The amount of characteristic elements (Al, Si, Cu and Fe) indicates the locations of the phases underneath the coatings and is beneficial to trace the evolution of coating formation and surface features on these phases respectively. Area 1 containing the highest amount of Si and lowest concentration of Al, Cu and Fe is corresponding to the coating formed on eutectic Si. Area 2 featured by the highest Al content among the three regions refers to the coating deposited on α-Al matrix. While the highest Cu and Fe quantities of area 3 indicate the coating on the typical overlapping region of Al₂Cu and β-Al₅FeSi intermetallics, which could hardly be distinguished after 480 s treatment.

The micro-pore size of the resultant coating on the overlapping of Al₂Cu and β-Al₅FeSi is 0.078-0.95 μm, 0.32-0.75 μm on α-Al matrix and 0.05-0.26 μm on eutectic Si respectively. Micro-cracks appear mainly in the coating grown on eutectic Si (Fig. 8(a)), which is confirmed by the clustering of Si (Fig. 8(b) and (e)). A high PBR value of 2.22 for SiO₂ indicates higher internal stresses of oxide mixtures mainly containing SiO₂. With smallest micro-pores on eutectic Si, the inefficient release of internal stresses consequently results in the generation of micro-cracks at these regions. Besides, the additional mismatch in an oxide mixture is prone to cause more micro-cracks in the coating as well.

3.2.4. Cross-sectional morphologies and element composition

With the goal of investigating the evolution of PEO coating, the cross-sectional morphologies in Fig. 9 demonstrates different growth stages of the coating (from 60 s to 480 s). It is obvious that the growth of coating on the different phases is dependent on the treatment time. Even after 60 s treatment, the oxide layer on the specimen is hard to recognize (Fig. 9(a)). However, the cross-sectional mapping result at 60 s (Fig. 9(a)) reveals a thin and discontinuous coverage of O on the surface of the specimen, chiefly associated with the disruption of the oxide layer by stable eutectic Si. It is worth noting that there is most likely an oxide layer already after 15 s and 30 s, but the metallographic preparation technique leads to rounding effects at the edges removing the thin film. Afterwards, protuberant porous oxides already start to deposit on the intermetallics of Al₂Cu and β-Al₅FeSi within 120 s, as shown by the distribution of Cu, Fe and O in Fig. 9(b), whereas only a thinner oxide layer covers the surface of eutectic Si.

A dense and continuous PEO coating coverage is firstly observed at 240 s (Fig. 9(c)). The coating on overlapping region of Al₂Cu and β-Al₅FeSi intermetallics is still thicker, mainly due to the lower melting temperature of the mixed oxides (Cu and Fe oxides) allowing faster sintering and coating growth. The gap between coating and substrate discloses a pore band in the interface. After that, the coating develops further up to 480 s. The coating becomes denser and more uniform in both morphology and thickness on all phases as presented in Fig. 9(d). This is most likely the results of the successive discharge breakdown in the weakest areas of the coating and deposition of reinforced coatings in these regions. The higher position of the interface of eutectic Si/coating relative to surrounding α-Al matrix/coating (Fig. 9(d)) indicates a delay of dissolution and oxidation of Si and a thinner oxide layer on eutectic Si. Simultaneously, the distribution of O on eutectic Si also suggests the highest stability of Si among all phases in the substrate.

For PEO coatings after 240 s and 480 s treatments, Fig. 10 reveals the coating evolution on the different phases in more details. The coating appears as undulant single-layer morphology on all phases at 240 s. The layer on eutectic Si is thin and jagged (Fig. 10(a)), while the layers on Al₂Cu and β-Al₅FeSi (Fig. 10(b) and (c)) are quite thick and dense, in agreement with results in literature [36]. The most visible pore bands arise up in the coating on the α-Al, adjacent to Al₂Cu and β-Al₅FeSi intermetallics, primarily attributed to the severe plasma discharges in these regions [37]. According to the element composition of the coating on different phases at 240 s (Fig. 10(g)), there seems to be less mixing of coating above eutectic Si. Si (SiO₂) still dominates the coating on eutectic Si and only little phosphate is incorporated. However, some Al₂O₃ seems to be present as well.

The coating thickens evidently on all phases after 480 s treatment (Fig. 10(d)-(f)). The average thickness values of the coating on α-Al matrix, Al₂Cu, β-Al₅FeSi and eutectic Si, measured from the cross-sections, are 2.62 ± 0.28 μm, 3.01 ± 0.29 μm, 3.06 ± 0.29 μm and 1.96 ± 0.38 μm respectively. A dense layer with smooth and sealed surface is successfully deposited on eutectic Si, despite the coating thickness remains the thinnest with respect to the other phases. The smooth conversion of the coating surface demonstrates that sufficient molten oxide was involved in coating formation at the final voltage of 350 V, even for the formation of SiO₂ (T_m>1600 °C).

Fig. 10(h) shows the chemical composition of the coatings on the four different phases at 480 s and a general increase of O content on eutectic Si from 240 s (less than 20 at.%) to 480 s (more than 25 at.%) can be seen. However, this seems to be more an artefact since there is more volume of the matrix region analyzed when the layer is thinner. Interestingly, the structure of the coating converts from a single-layer morphology to double-layer morphology on the matrix, Al₂Cu and β-Al₅FeSi intermetallics as shown in Fig. 10(e) and (f) after 480 s. It may arise from a change of main growing direction of the coating at 240 s. The subsistent thick coating impedes the breakdown discharges, which further impairs the outwards growing of the coating while inwards growth rate of the coating increases [3]. With increase of treatment time, an outer layer with visible micro-pores and an inner compact layer free of micro-pores develop eventually.

At the very beginning, the boundary of intermetallics Al₂Cu is initially dissolved (Fig. 4(a)) with the assistance of applied electrical field. Prior to anodizing, a native oxide layer formed in air already covers the alloy (Fig. 11(a)). The oxide layer on Al₂Cu intermetallics is not as protective as that on α-Al matrix as a result of the high content of Cu. Inevitably, the intermetallics underneath are exposed directly to the electrolyte and especially at the interface of Al₂Cu/α-Al the native oxide layer compromises even more defects [38]. Consequently, the electrical field-assisted dissolution is preferentially initiated and enhanced at the verge of Al₂Cu intermetallics. Afterwards, conversion products start to deposit in these specific regions by the combination of Al³⁺/Cu²⁺ from dissolved matrix and PO₄^3-/OH- originating from the electrolyte. They form of low soluble Al/Cu-based oxides and phosphates (Figs. 5(a) and (b) and 11 (b)). The enrichment of Fe will produce residual defects in the native oxide layer on β-Al₅FeSi providing a path for the penetration of electrolyte. The presence of Si in β-Al₅FeSi tend to change the current flow into less resistive α-Al matrix around [39] and reduce the anodizing efficiency. Overall, β-Al₅FeSi presents a lower oxide growth rate relative to intermetallic Al₂Cu (Fig. 5). With the increased voltage, conversion products gradually spread to the surface of intermetallics and accumulate further (Fig. 11(b)).

The unexpected peeling-off (Fig. 4(d)) interrupts the collection of conversion products after the breakdown voltage is reached. It could be related to the thermal and mechanical stresses generated by the first discharges. Due to the excess of flaws within high Cu contained conversion products the poor adhesion to the substrate beneath is also a possible factor (Fig. 11(c)). The disappearance of protuberant conversion products potentially provides more weak regions for the plasma breakdown discharges and the coating growth will be enhanced by the intensive and frequent discharges [40].

Plasma discharges after their extinguishing, are responsible for the typical surface morphology with micro-pores on the coating surface. When the increased voltage reaches the critical breakdown voltage, the breakdown of dielectric layer takes place hence leaves behind the characteristic micro-pores near the coating surface and micro-channels within the coating [41].

The micro-pores from initial discharges are largely distributed on the top of α-Al matrix (Fig. 6), demonstrating the ease of breakdown discharge in corresponding barrier layers. Interestingly, these micro-pores are preferentially surrounding Al₂Cu intermetallics (Fig. 6(a)). A possible explanation is associated with the modified structure and composition of the preferential barrier layer on α-Al matrix next to the interface of Al₂Cu/α-Al.

Severe gas generation during anodizing at the layer/electrolyte interface prevents the penetration of electrolyte into the layer, insulating the substrate and creating numerous resembling capacitors in the interface barrier layer/electrolyte. Once the voltage exceeds the breakdown voltage, gas discharges take place accompanying with the breakdown of the barrier layer. Considering the layers having excess of defects, i.e. the layers on Al₂Cu containing high content of copper oxide [42], are more prone to pass through by the current flow and sustain dissolution and oxidation of the intermetallics underneath rather than to be a barrier element of the capacitors before reaching a specific thickness. In regards to layers of bare imperfection, i.e. the crystalline layer on α-Al, although served as a part of the capacitors, are hardly to be broken down due to the low dielectric permittivity of Al₂O₃ (7.5-15) relative to CuO (around 18) [43]. Nevertheless, layers around the interface of Al₂Cu/α-Al are modified either by morphology or composition (Fig. 5(b)) arising from diffusion of Al³⁺/Cu²⁺ in these regions compared to layers strictly formed on Al₂Cu and α-Al. Subsequently, layers at the interface of Al₂Cu/α-Al mixed with Al₂O₃ offering good passivity and CuO with high dielectric permittivity is precisely subjected to the breakdown.

The continuous increase of voltage determines the growth of PEO coatings. The cycles of arcing, melting, sintering and quenching of the coating during pulsed PEO processing cause the sintered appearance of the coating on the overlapping regions of Al₂Cu and β-Al₅FeSi (Fig. 7(d)). Actually, lower melting points of FeO (1377 °C), Fe₂O₃ (1565 ℃) and CuO (1450 °C) relative to higher melting points of Al₂O₃ (2072 °C) can be responsible for the more sintered morphologies with larger micro-pore size compared to the denser coatings on the surrounding α-Al matrix.

The continuous PEO coatings possess the smallest micro-pores on eutectic Si (Figs. 7, 8 and 11(d)), determined by less intensive breakdown discharges on eutectic Si. At this stage, all phases are on the same potential. In terms of the low conductivity of eutectic Si, less current can flow through the combination of Si/SiO₂ (semiconductor/oxide) than through the combination Al/Al₂O₃. The wider band gap of SiO₂ compared with other oxides also contributes to the lower current passing. Weaker current flow makes the ignition of energetic discharges on eutectic Si more difficult as well as being responsible for lower intensity of discharges. Besides this, the hardship in breaking down the SiO₂ barrier layers with low dielectric permittivity 3.8 on eutectic Si [43] further prevent the formation of large-sized micro-pores. Generally, every single discharge is a transient process to generate energetic species from the collisions between high-energy electrons with gas molecules and atoms in gasification. These energetic species then react with the barrier layers by breaking the chemical bonds in barrier layers, revealing a typical breakdown process [44]. Highest ionization energy of SiO₂ (11.43 eV) relative to Al₂O₃ (10.5 eV), CuO (9.41 eV) and FeO (8.56 eV) [45] eventually moderates strength of reactions on the barrier layer of eutectic Si.

With the thickening of the coating, the breakdown of the coating becomes more difficult and the outward growth of the coating reaches a limit soon. In this case, the bottom of some large-sized discharge channels turn into new reaction zones for the inward growth of the coating. The opposite directions in coating growth finally give rise to the double layer morphology in the coating as shown in Fig. 10(e) and (f) and depicted in Fig. 11(e). At this stage, gas discharges are prone to take place at the bottom of micro-pores [46], from which the pore bands generate. The high electric field forces the ionized gases inward from the electrolyte/coating interface to the substrate and drives the aluminum ions outward via the discharge channels. The inward ionized gases, filling the discharge channels at high pressure, consequently constrain the accesses of electrolyte into the discharge channels. A quenching effect of the electrolyte on the solidification of the newly formed molten oxides is hence prevented. As a result, a micro-pore free inner layer is deposited next to the substrate.

(1) Al₂Cu intermetallics, β-Al₅FeSi and eutectic Si in AlSi9Cu3 alloy exhibit significant influence on phase formation and coating growth during PEO processing mainly attributed to their different chemical-, electrochemical- and physical properties. The coatings formed in phosphate electrolyte are mainly composed of γ-Al₂O₃, mullite (3Al₂O₃·2SiO₂) and an amorphous phase.

(2) The involvement of the phases in the PEO process (discharge formation) after initial conventional anodic oxidation is α-Al matrix, followed by Al₂Cu and β-Al₅FeSi and finally eutectic Si in sequential order. The majority of initial micro-pores (discharges) appear preferentially on the α-Al matrix surrounding Al₂Cu intermetallics, indicating more active sites for plasma discharges at these regions.

(3) Resultant PEO coatings show distinct surface morphologies on different phases. The coating on Al₂Cu and β-Al₅FeSi intermetallics reveals a sintered morphology with large micro-pores. The thinnest coating with smallest micro-pores and some micro-cracks is found on eutectic Si.

(4) A double-layer morphology in the resultant coating is observed on all phases except of eutectic Si, characterized by a porous outer layer and a compact inner layer, due to the reverse growing directions (outward firstly then inward) of the coating.

(5) The alloy AlSi9Cu3 is a good model alloy for PEO treatment of multi-material mix, revealing the problems of delayed sparking and different coating growth/phase composition if two different materials should be treated.