Microstructure and Microhardness of High-Strength Aluminium Alloy Prepared Using High-Speed Laser Fabrication

Abstract

As a recently developed high-strength aluminium alloy used specifically for laser additive manufacturing, AlMgMnSc alloy possesses superior mechanical properties and excellent processability. Extreme high-speed laser deposition (EHLD) is a novel surface-modification technique, which is characterised by high depositing speed, rapid cooling, rate and minimal dilution rate. To offer a new method for surface repairing high-strength aluminium alloys, an AlMgMnSc alloy coating, containing two deposition layers, is prepared on a 6061 aluminium-alloy axle using the EHLD technique. Meanwhile, the microstructure, composition distribution, and microhardness variation of the fabricated coating are studied. The results reveal that the coating is dense and crack-free, which is well-bonded with the substrate. Additionally, layer 1 is mainly composed of large columnar and equiaxed grains, while layer 2 consists of a fully equiaxed grain structure with an average grain size of about 4.5 μm. Moreover, the microhardness of the coating (about 104~118 HV) is similar to the substrate (about 105 HV), proving the feasibility of repairing high-strength aluminium alloys using AlMgMnSc alloy powders through the EHLD technique.

1. Introduction

Extreme high-speed laser deposition (EHLD) is a novel material-surface-modification technique, which was developed by Fraunhofer Laser Technology Research Institute in 2013 [1,2]. In conventional laser deposition, primary energy is deposited into the substrate instead of into powders. As for the EHLD, most of the laser energy is absorbed by powders, decreasing the heat input to the substrate remarkably [3]. During the EHLD process, raw materials are fed by a coaxial powder feed nozzle, with the focus of the powder flow above the substrate. Most of the laser energy is absorbed by the powders before they reach the substrate; only a small melting pool is formed on the substrate surface by residual laser energy. As a result, powder particles are melted by the laser beam before hitting the substrate, entering the melting pool in a molten or semi-molten state [4,5]. Consequently, the EHLD has the advantages of high depositing speed, rapid cooling rate, and minimal dilution rate, providing high economic benefits and satisfactory material performance [2,6]. As is reported, the processing speed and deposition efficiency of the EHLD can be up to 500 m/min and 96%, respectively [7].

Currently, the EHLD is mainly used to modify the surfaces of Ni-based alloy, Fe-based based alloy, amorphous alloys, high-entropy alloys, and metal matrix composite coatings [6,8,9,10,11,12,13,14,15,16]. Yuan et al. [1] deposited Ni45 alloy coatings on 45 medium carbon steel substrates by a traditional laser-deposition technique and the EHLD, respectively. It is reported that the coating fabricated using the EHLD has a lower roughness and a smaller thickness when compared with the traditional laser-deposition technique. Meanwhile, variations in processing parameters lead to changes in the solidification conditions of the melting pool. A fine equiaxial dendritic structure could be obtained under a high laser cladding rate. Zhang et al. [4] established a powder flow model by using a transient bidirectional coupling calculation method to investigate the effects of laser power, carrier-gas flow rate, and shielding-gas flow rate on particle temperature profiles during the EHLD process. According to their results, the carrier-gas and protective-gas flow rates are key factors for determining the maximum elevation angle of the powder flow. With the increase in carrier-gas flow, there is an increase in the maximum inclination angle of the powder. On the other hand, the increase in shielding-gas flow could lead to a decrease in the maximum inclination angle of the powder. Koß et al. [7] investigated the maximum surface rate and deposition rate by depositing an AlSi10Mg coating on an AlMgSi0.5 alloy substrate using the EHLD with a conventional 8 kW IR-laser beam source. The results demonstrate the feasibility of fabricating an aluminium coating on an aluminium-alloy substrate through the EHLA. At a processing speed of 100 m/min, a maximum surface rate of 800 cm²/min and deposition rate of about 22.5 g/min could be obtained. Meanwhile, a maximum surface rate of 1300 cm²/min and a deposition rate of about 34 g/min were achieved at a processing speed of 200 m/min. It is proposed that it is possible to increase the surface and deposition rate by increasing the laser power using the EHLD.

High-strength aluminium alloys are used in a broad range of industries, such as the automobile, aerospace, and construction industries and mechanical and electrical engineering [17,18,19]. The world’s top-three largest consumers of primary aluminium are China (33.3 million tons), the USA (4.6 million tons), and Germany (2.1 million tons) [7]. As is reported, aluminium alloys make up 60–80% of the weight of a commercial aircraft and more than 50% of the structural quality of rockets and aircraft [20,21]. The 6xxx series alloy is a kind of high-strength aluminium alloy. It is reported that the application proportion of the 6xxx series alloy in civil aircraft is about 6% [20]. A typical 6xxx series alloy is 6061 aluminium alloy, which is a heat-treated, precipitation hardened, alloy and includes magnesium (Mg), silicon (Si), and iron (Fe) as its major alloying elements [22]. The 6061 aluminium alloy offers low density, high strength-to-weight ratio, good ductility, excellent formability and weldability, good corrosion resistance, and low cost [22]. Currently, it is mainly adopted to fabricate aircraft wings, fuselages, and fuel tanks and in wheel rims, wheel spacers, truck bodies, car frames, bicycle frames, bike frames, container bodies, goods wagons, carriages, trams, driveshafts, valves, etc. [22,23].

The long-term performance of high-strength aluminium-alloy components under certain conditions is limited due to their high coefficient of friction, poor wear characteristics, and low hardness. Additionally, it is reported that about 423.2 kWh of electricity and about 558 Nm³ of natural gas are needed to produce 1 t of primary aluminium, while recycling requires only about 5% of the energy [7]. Consequently, surface repair and modification are of great importance for high-strength aluminium alloys to reduce material quantity and manufacturing costs.

As is known, the preparation of traditional high-strength aluminium alloys through laser processing is challenging because they normally present high susceptibility to cracking due to large liquidus-to-solidus temperature ranges and high solidification rates during fabrication [24,25,26,27,28]. Tan et al. [25] analysed the influence of processing parameters on the cracking of a 2024 aluminium alloy prepared by laser processing. It is suggested that a reduction in the fraction of large elongated columnar grains could lead to a decrease in hot cracks. However, it is impossible to eliminate columnar grains by adjusting processing parameters. Thus, hot cracks in laser-processed 2024 aluminium alloys could hardly be completely avoided. Wojciech et al. [26] optimised the processing parameters of a 7075 aluminium alloy fabricated through laser processing. Their results indicate that solidification cracks along columnar grain boundaries could hardly be totally eliminated. Furthermore, intergranular micro-segregation of Zn, Mg, Cu, and Si are observed. It is considered that Si may result in a high cracking susceptibility by increasing the stability of the liquid film. Yao et al. [28] studied the formation mechanism of cracks in a laser-processed 2024 aluminium alloy. They proposed that the transverse strain rate of columnar grains could easily exceed the sum of the transverse expansion rate of grains and the filling rate of liquids at the grain boundaries due to high Cu and Mg contents. In addition, the high cooling rate and heat input involved in laser processing reduce the filling rate of liquids at grain boundaries. As a result, a 2024 aluminium alloy presents strong thermal crack sensitivity during laser processing. Therefore, the application of laser processing in surface repair and modification of high-strength aluminium alloys is significantly limited.

Very recently, Sc-containing high-strength aluminium alloys have been designed specifically for laser additive manufacturing (LAM) [24,29,30,31,32,33,34,35,36]. In LAMed AlMgMnSc alloys, primary Al₃Sc particles would precipitate in the melting pool before solidification of the α-Al, providing heterogeneous nucleation sites for α-Al grains [30]. As a result, the formation of equiaxed fine grains is promoted, which mitigates solidification cracking significantly. Meanwhile, the addition of Sc remarkably improves the mechanical properties of aluminium alloys. Wang et al. [31] studied the influence of solidification conditions on the microstructure and mechanical property of an aluminium alloy fabricated through laser-directed energy deposition. It was found that an equiaxed grain structure could be obtained at a low cooling rate, while a heterogeneous grain structure containing a fine grain band and a fan-shaped coarse grain was formed at a high cooling rate. After ageing, the yield strength of the sample prepared at a high cooling rate is enhanced to ~2 times that of the sample prepared at a low cooling rate. Hua et al. [32] prepared aluminium alloys with a fully equiaxed grain structure with grain sizes of 2–30 μm through laser-directed energy deposition. Most of the primary Al₃Sc particles were found along grain boundaries. The deposited alloy presents a certain but not obvious anisotropy, since the texture index and strength are only slightly larger than 1. Yield strength, tensile strength, and elongation of the specimen in the horizontal direction and the building direction are 399.87 MPa, 220.96 MPa, 9.13% and 385.40 MPa, 219.40 MPa, 8.24%, respectively.

Although the successful preparation of AlMgMnSc alloys through LAM suggests that it could be possible to modify the surfaces of aluminium alloys through the EHLD technique, research in this field is seldom conducted. The microstructure characteristics and mechanical behaviour of the EHLDed AlMgMnSc alloy coating remain unclear. Therefore, the EHLD approach was used in this work to fabricate an AlMgMnSc alloy coating on a 6061 aluminium-alloy axle. Additionally, the microstructure, composition, distribution, and microhardness variation of the prepared coating were analysed. Meanwhile, the microstructure formation mechanism of the deposited coating was investigated. The aim of this paper is to offer a comprehensive understanding of the microstructure and mechanical property of the EHLDed AlMgMnSc alloy coating, providing theoretical guidance for surface repairing of a high-strength aluminium alloy using the EHLD technique.

2. Materials and Methods

An annealed wrought 6061 aluminium-alloy axle with a diameter of 60 mm was applied as a substrate. The composition and microhardness of the substrate are about Al-1.12Mg-0.18Cu-0.12Mn-0.17Zn-0.19Cr-0.06Ti-0.61Si-0.55Fe (wt.%) and 105 HV, respectively. Spherical AlMgMnSc alloy powders with particle sizes ranging from 15 to 53 μm were adopted as feedstock material (D₁₀: 17.30 μm, D₅₀: 33.15 μm, and D₉₀: 55.58 μm). The apparent density and tap density of the powders are 1.40 g/cm³ and 1.7 g/cm³, respectively. Figure 1 illustrates the SEM images of the powders. As shown in Figure 1a, the powders have a spherical shape. It is also noticed that some planetary powders are distributed on the surfaces of powders (see Figure 1b). The nominal composition of the AlMgMnSc alloy is Al-5Mg-0.4Sc-0.35Zr-0.9Mn-Si-0.6Ti-0.6Cu-0.25Cr (wt.%). The deposition process was carried out using an LDF3000-40 EHLD system (Laserline GmbH, Muelheim, Germany), which is composed of a high-speed lathe and laser and powder feeding system. The laser adopted in this paper is a LDF3000-40 semiconductor pulse laser (Laserline GmbH, Muelheim, Germany). The pulse duration and wavelength of the laser are 100 ms and 960~1060 ± 10 nm, respectively. During processing, the laser beam is focused by a focusing lens group. A detailed introduction to the EHLD system can be found in Ref. [1]. In this paper, the coating was prepared by depositing two AlMgMnSc alloy layers using the following processing parameters: a powder feed rate of 6 g/min, laser spot diameter of 1.2 mm, laser power of 1500 W, scanning speed of 5.5 m/min, and overlapping ratio of 50%. According to our previous work on the optimization of processing parameters, crack-free EHLDed AlMgMnSc alloy coatings with a metallurgical bond to the substrate can be obtained by using the above processing parameters. Before deposition, the surface of the 6061 aluminium-alloy axle was polished and cleaned with alcohol and acetone. During the deposition process, argon was adopted as a powder carrier gas and shielding gas. A schematic illustration of the EHLD process is shown in Figure 2.

After deposition, the fabricated sample was sectioned into pieces using electrical discharge wire cutting to analyse the microstructure, composition variation, and microhardness distribution of the longitudinal section of the prepared coating. The samples were ground with SiC grit paper and then polished using the standard mechanical polishing method. Mechanical polishing was conducted by adopting a Tegramin-20 polishing machine (Struers, Copenhagen, Denmark), and the duration of mechanical polishing is about 1 min. After mechanical polishing, a VibroMet 2 vibratory polisher (BUEHLER, Lake Bluff, IL, USA) was adopted to polish the surfaces of the samples. During vibratory polishing, SiO₂ turbid liquid was utilised as a grinding polishing media and the polishing time was 4 h. Microstructure, crystallographic orientations, grain size distributions, and composition variation were determined by a JSM7900F scanning electron microscope (SEM) (JEOL, Tokyo, Japan) equipped with energy-dispersive X-ray analysis (EDX). Microhardness distributions along depth direction were tested by using an FM-800 Microhardness tester (Future-Tech, Tokyo, Japan), with 10 s loading time and under loads of 50 g. The average hardness of each position was calculated from three measurements.

3. Results and Discussion

To reveal the microstructure formation mechanism of the prepared AlMgMnSc alloy coating, the change of composition, grain morphology, grain size, and crystallographic orientation along the deposition direction are analysed. In addition, to verify the feasibility of repairing aluminium alloys by utilising the EHLD technique and AlMgMnSc alloy powders, the microhardness distribution along the depth direction is measured.

Variation in composition along the depth direction of the fabricated sample was measured by EDX, and the result is given in Figure 3. As is shown, with an increase in the distance away from the substrate, the contents of Fe and Zn decrease while the contents of Mg, Sc, Zr, Si, Mn, Ti and Cr are increased. The average compositions of layers 1 and 2 are Al-3.5Mg-0.36Sc-0.16Zr-Si-0.5Ti-0.22Cr-0.8Mn-0.17Fe-0.6Cu-0.2Zn (wt.%) and Al-4.4Mg-0.42Sc-0.32Zr-Si-0.55Ti-0.24Cr-0.9Mn-0.15Fe-0.6Cu-0.01Zn (wt.%), respectively. During deposition of layer 1, part of the substrate is melted and mixed with newly deposited AlMgMnSc alloy. Thus, a layer with a composition between 6061 aluminium alloy and AlMgMnSc alloy is established. As deposition continues, the composition of layer 2 is more close to the AlMgMnSc alloy. Finally, a coating with changed compositions is formed.

Figure 4 shows SEM images of a cross section of the fabricated sample. An overview of the coating in Figure 4a indicates that the coating is dense and crack-free. Also, pores with an average diameter smaller than 100 μm are observed. The boundary between the substrate and the coating is illustrated in Figure 4b, revealing that the interface between the coating and the substrate is smooth, continuous, and well-bonded. The microstructure of layers 1 and 2 in Figure 4c,d indicates that both layers are composed of precipitation phases and an α-Al matrix. The microstructure of layer 1 in Figure 4c reveals that most of the precipitation phases exhibit a lath-shaped morphology. Also, a small amount of nano-scaled precipitation phases with a dot-shaped morphology are visible. Meanwhile, the precipitation phases in layer 1 mainly precipitate around the grain boundaries of α-Al, displaying a lamellar structure. The microstructure of layer 2 in Figure 4d shows that nano-scaled precipitation phases with particle-like morphology disperse evenly through equiaxed α-Al grains.

To further analyse the microstructural characteristics of different layers, electron backscattered diffraction (EBSD) analysis is adopted. The EBSD results of layer 1 are shown in Figure 5. The orientation map in Figure 5a indicates that layer 1 has a unique heterogeneous grain structure, which can essentially be divided into three regimes based on the variation in grain size and morphology: (i) equiaxed grains at the bottoms of melting pools (EGB), (ii) columnar grains grown from the bottoms of melting pools (CGB), and (iii) equiaxed grains at the tops of melting pools (EGT). The EGB mainly distributes at the melting boundary close to the previously deposited layer, while the CGB primarily distributes at the melting boundary adjacent to the substrate. In addition, columnar grains preferred to grow along the direction perpendicular to the melting boundary. From Figure 5b, it is noticed that the distribution of grain size in the EGB region is not uniform. In the colony of fine equiaxed grains (FEGB), there is a band of relatively coarse equiaxed grains (CEGB).

To further analyse the difference between crystallographic orientations and the grain size of different regimes in layer 1, the texture index and grain size distributions are quantitatively compared in Figure 6. Pole figures (PFs) of the FEGB, CEGB, and EGT regions in Figure 6a–c illustrate that equiaxed grains in layer 1 are randomly oriented, which can be confirmed by the random colour in Figure 5. The maximal texture indices of the FEGB, CEGB, and EGT regions are 1.936, 4.314, and 5.774, respectively. The inverse pole figure (IPF) of the CG region in Figure 6d suggests that columnar grains mainly grow along the crystallographic direction of [212] and [103]. Grain-size data calculated based on the EBSD spectrum in Figure 6e–g show that the grain size distributions of the FEGB, CEGB, and EGT regions are 0.2~1.8 μm, 0.2~2.8 μm, and 1.7~41.9 μm, respectively. The average grain sizes of the equiaxed grains in the FEGB, CEGB, and EGT regions are about 1.0 μm, 1.9 μm, and 16.0 μm, respectively.

The grain structure of the AMed AlMgMnSc aluminium alloy is primarily determined by the solidification condition of the melting pool and precipitation of primary Al₃(Sc, Zr) particles [37]. During solidification, a high temperature gradient (G) at the solidification interface and low solidification velocity (V) favour the formation of columnar grains growing along the direction parallel to the direction of the maximum thermal gradient. By contrast, high V and low G promote the formation of equiaxed grains [38,39]. Additionally, primary Al₃(Sc, Zr) particles tend to precipitate from liquid metals before the onset of α-Al solidification. The similar lattice constants of α-Al and Al₃(Sc, Zr) make the Al₃(Sc, Zr) particles effective nucleation cores for the following solidification of α-Al grains [37]. As a result, the formation of a fine equiaxed grain structure is strongly promoted.

During the deposition of layer 1, the G and cooling rate (R) of the melting pool are relatively large, due to the low temperature of substrate. Rapid cooling during solidification may suppress precipitation of primary Al₃(Sc, Zr) particles [31]. Additionally, G usually enhances with the increase of depth in the melting pool, while V is the highest at the top of the melting pool [40,41]. Therefore, columnar grains prefer to form at the bottom of the melting pool, while equiaxed grains are obtained at the top of the melting pool [42]. Consequently, layer 1 mainly contains large columnar and equiaxed grains. The addition of Sc does not result in a widespread presence of fine equiaxed grains. In layer 1, it is also noticed that only a small amount of fine equiaxed grains are formed at the melting boundary close to the previously deposited layer. The existence of fine equiaxed grains could be attributed to the high temperature of the previously prepared layer, which leads to a decrease in G and R. Under a relatively low R, primary Al₃(Sc, Zr) particles would be obtained in the melting pool, providing heterogeneous nucleation sites. Simultaneously, a low G promotes the formation of equiaxed grains. Generally, variation of local solidification conditions results in heterogeneity in the grain structure of layer 1.

EBSD results of layer 2 in Figure 7 show that the grain structure of layer 2 is quite different from that of layer 1, indicating that these two layers experience different modes of solidification. Orientation maps in Figure 7a–d reveal that no columnar grain is formed in layer 2, and most grains exist in the form of equiaxed grains. The fully equiaxed grain structure in layer 2 is similar to the AlMgSc aluminium alloys prepared using a laser-directed energy deposition technique [31,32]. Despite the similarity in grain morphology, the distribution of grain size in layer 2 is not uniform. According to the variation in grain size, layer 2 can be divided into three regimes: (i) fine equiaxed grains at the bottoms of melting pools (FEG), (ii) equiaxed grains with medium size at the centres of melting pools (MEG), and (iii) coarse equiaxed grains at tops of melting pools (CEG).

To study the differences in crystallographic orientations and grain size of different regimes in layer 2, the texture index and grain size distributions are quantitatively investigated in Figure 8. PFs of the FEG, MEG, and CEG regions in Figure 8a–c show that layer 2 presents strong anisotropy, and grains in layer 2 have no obvious preferred orientation. The maximal texture indices of the FEG, MEG, and CEG regions are 2.054, 1.465, and 2.149, respectively. Grain-size data obtained from the EBSD spectrum in Figure 8d–f show that the grain size distributions in the FEG, MEG, and CEG regions are determined to be 0.2~1.8 μm, 0.7~5.0 μm, and 0.7~12.1 μm, respectively. The average grain sizes of the equiaxed grains in the FEG, MEG, and CEG regions are about 1.2 μm, 2.2 μm, and 5.6 μm, respectively. Layer 2 displays an average grain size of about 4.5 μm. Compared with layer 1, it is clear that the distribution of grain size in layer 2 is more uniform. According to previous research, a fully equiaxed grain structure with grain sizes of 2–30 μm could be obtained by using laser-directed energy deposition [32]. As can be seen, the alloy prepared through the EHLD in this work is smaller than the laser-directed energy-deposited alloy.

During the deposition of layer 2, G and R are relatively small due to the high temperature of layer 1. Compared with layer 1, precipitation of primary Al₃(Sc, Zr) particles is more homogeneous in the entire melting pool. The enrichment of primary Al₃(Sc, Zr) particles leads to the formation of a fully equiaxed grain structure. The formation process of fine equiaxed grains is illustrated in Figure 9. In addition, Figure 7 indicates that the distribution of grain size in layer 2 is not uniform. This may be ascribed to the convection of liquid in the melting pool [43]. It is speculated that liquid flow leads to a non-uniform distribution of primary Al₃(Sc, Zr) particles, resulting in the variation in grain size.

As is reported by previous works, Sc-containing aluminium alloy fabricated by selective laser melting normally presents a bimodal grain structure. That is, a band of equiaxed grains with no preferential crystallographic orientation forms at the melting boundary, while columnar grains growing along the deposition direction dominate the centre of the melting pool [44,45]. On the other hand, a fully equiaxed fine grain structure can be obtained in Sc-containing aluminium alloy prepared through laser-directed energy deposition, due to a relatively low cooling rate and temperature during deposition [32]. As for the prepared coating in this paper, the unique fabrication process of the EHLD leads to different microstructure features. The variation in solidification conditions results in a changed grain morphology along the building direction. That is, layer 1 displays a heterogeneous grain structure, while layer 2 consists of a fully equiaxed grain structure.

Figure 10 presents variations in microhardness along the depth direction of the fabricated sample and micrographs of microhardness indent marks. As can be seen, the microhardness of layers 1 (about 104~112 HV) and 2 (about 107~118 HV) are similar to the substrate (about 105 HV). Hence, it could be suggested that using AlMgMnSc alloy powders as raw materials to repair the surfaces of traditional high-strength aluminium-alloy components via the EHLD technique is applicable. It is also noticed that microhardness fluctuates obviously within the deposited layers. As illustrated in Figure 5 and Figure 7, the fabricated coating presents a mixed microstructure. That is, there is a change in grain morphology and grain size along the deposition direction. Therefore, it is speculated that the fluctuation in microhardness may be ascribed to the uneven distribution of the microstructure.

To sum up, this paper reveals the microstructure and microhardness of the EHLDed AlMgMnSc alloy coating for the first time. It is found that a change in the solidification conditions within the melting pool results in the variation in grain morphology along the deposition direction. Moreover, fully equiaxed fine grains can be obtained with an increase in the deposition layer. In addition, the fabricated coating exhibits a similar microhardness to the 6061 aluminium-alloy substrate. This paper offers a deep insight into the microstructure and mechanical performance of the EHLDed AlMgMnSc alloy coating.

4. Conclusions

In this paper, an AlMgMnSc high-strength aluminium-alloy coating was prepared on a 6061 aluminium-alloy axle using the EHLD technique. The grain structure, composition distribution, and microhardness variation of the fabricated coating were studied.

The results reveal that the prepared coating is dense and crack free and is well-bonded with the substrate. Additionally, with an increase in distance away from the substrate, there is a change in chemical compositions along the deposition direction. Moreover, due to the variation of local solidification conditions, layer 1 mainly contains large columnar and equiaxed grains. Only a small amount of fine equiaxed grains are formed at the melting boundary. The average grain size of the equiaxed grains in the FEGB, CEGB, and EGT regions in layer 1 are about 1.0 μm, 1.9 μm, and 16.0 μm, respectively. Meanwhile, layer 2 consists of a fully equiaxed grain structure with no preferential crystallographic orientation. The distribution of grain size in layer 2 is not uniform. The average grain sizes of the equiaxed grains in the FEG, MEG, and CEG regions in layer 2 are about 1.2 μm, 2.2 μm, and 5.6 μm, respectively. Furthermore, the microhardness of the coating (about 104~118 HV) is similar to the substrate (about 105 HV), proving the feasibility of repairing the surfaces of traditional high-strength aluminium-alloy components using AlMgMnSc alloy powders through the EHLD technique. Compared with conventional laser deposition, the EHLD technique presents higher depositing efficiency. In the meantime, primary Al₃(Sc, Zr) particles in AlMgMnSc alloy promote the formation of equiaxed fine grains, inhibiting the generation of solidification cracking remarkably. Also, the addition of Sc significantly improves the mechanical properties of the alloy. By combining the advantages of the EHLD technique and an AlMgMnSc alloy, this paper provides a new approach to rapidly modify the surfaces of aluminium-alloy components with high performance.