Effect of melting modes on microstructure and tribological properties of selective laser melted AlSi10Mg alloy

ABSTRACT

This paper focuses on the effect of melting modes on microstructural evolution and tribological properties of AlSi10Mg alloy fabricated by selective laser melting (SLM). The results showed that the microstructures of SLM AlSi10Mg consisted of primary α-Al surrounded by cellular Si networks (∼500 nm) when fabricated in conduction mode, but has a finer cellular-like Si phase (∼200 nm) when fabricated in keyhole mode. The strong convection caused by the melt reflow and Marangoni convection under keyhole mode also resulted in deposition of nano-scale Si particles at the bottom of the molten pool. The SLM AlSi10Mg fabricated in keyhole mode exhibited better wear resistance than that fabricated in conduction mode. Compared to traditional as-cast specimens, both SLM specimens showed better wear resistance due to the unique cellular-like networks. The SLM technique offers a new approach for material processing that can be used to refine microstructures for improved tribological properties.

KEYWORDS:

• Additive manufacturing
• AlSi10Mg
• melting mode
• tribological behaviour
• nanoindentation

1. Introduction

Many engineering parts are subject to wear during service, which not only reduces their service life, but also increases energy consumption and generates both noise and vibration. Thus, good wear performance is a critical parameter for many materials. AlSi10Mg is a Al–Si alloy that has several advantages, including being lightweight, good wear resistance, and excellent thermal conductivity (Aboulkhair et al. 2019). It is widely used in both automobiles and aircraft in engine blocks, pistons, and hinges (Wahab et al. 2019; Tang et al. 2020). With the ever-increasing demand for lightweight, structural and functional integration of aluminum alloy components, the production of complex-shaped, thin-walled, and high-precision parts has become the primary development trend at present (Demir and Biffi 2019; Leary et al. 2019). Traditional aluminum alloy forming processes have long production cycles and high production costs, which have difficulty meeting the growing current technical requirements.

Selective laser melting (SLM), as a typical metal additive manufacturing technology, can fabricate complex metal parts layer by layer directly using digital models. This technology has the ability to produce components with unlimited flexibility and high accuracy, thus, minimising both processing cycles and material waste (Zhang et al. 2019). Owing to the latest developments in laser technology, SLM equipment can melt and consolidate low laser absorption and high laser reflection materials, such as aluminum alloys and magnesium alloys (Li et al. 2018; Bi et al. 2020; Long et al. 2018; Yin et al. 2019). During laser melting, high volumetric energy density (E) at high laser power (P) and low scanning speed (v) can be used to change the melting mode from conduction mode (CM) to keyhole mode (KM), resulting in a transition of the melt shape from shallow and semicircular to deep and narrow (Qi et al. 2017). In KM, the high E not only vaporises material and forms a plasma, but also enhances laser absorption and drills a vapour capillary at the bottom of the molten pool (Dowden 2009). Thus, the KM in 3D printing can lead to excessive pore defects. Considerable research has focused on the keyhole formation mechanism and the keyhole threshold to avoid the occurrence of this phenomenon (Cunningham et al. 2019; Teng et al. 2017; Shrestha and Kevin Chou 2019; Thanki et al. 2019). However, Aboulkhair et al. (2016) pointed out that use of the KM can be acceptable for AlSi10Mg due to the formation of pore-free molten pool during depositing. Yang et al. (2016) found that the CM exhibits superior formability, while the KM produces a better combination of strength and ductility. Therefore, it is important to systematically investigate the microstructural evolution and related tribological properties of SLM components fabricated by the two different modes. It is difficult to evaluate the effect of microstructural evolution on wear performance using traditional wear tests. Recent works have been focused on nanoindentation tests to solve this problem (Okoro et al. 2019; Zhou, Xu, et al. 2019). For example, the relationship between the thickness of α′ and tribological properties of SLM Ti-13Nb-13Zr alloys can be well established through nanoindentation tests (Zhou, Yuan, et al. 2019).

The objective of this study was to explore the microstructures and tribological properties of SLM AlSi10Mg fabricated under different conditions, i.e. in CM and in KM, and to understand the related wear mechanisms. For comparison, wear tests were performed on an as-cast specimen to understand the influence of the microstructure on tribological properties.

2. Experimental

Gas-atomized AlSi10Mg powder used in this study was prepared using a close-coupled gas atomisation system under high-purity argon (99.99 wt.%). The composition (wt.%) of the powder was 8.21Si, 0.21Fe, 0.01Cr, 0.02Mn, 0.82Mg, 0.02V, 0.03Ni, 0.11O, with the balance Al. Two groups of specimens (10 mm × 10 mm × 10 mm) were fabricated using a FS271M SLM machine (Farsoon, Inc., China) equipped with a maximum power of 500 W Yb: YAG fibre laser. To obtain the specimens under two melting modes, the SLM parameters employed were: P = 175 W and v = 694 mm/s; or P = 250 W and v = 500 mm/s. The laser beam diameter (d), hatching spacing (h) and layer thickness (t) were kept constant at 90, 120 and 50 μm, respectively. All the processing was performed under an inert and high purity Ar gas containing no more than 100 ppm oxygen. The scanning strategy plays a crucial role in forming quality. Short scan vectors could decrease the residual stress, and the proper rotation angle could reduce part deformation without detriment to strength or part porosity (Robinson et al. 2019). Thus, a ‘checkerboard’ scanning strategy with 67° rotation between each layer was utilised for the bulk volume. It is well known that compressive residual stress has a positive effect on wear resistance. The tensile stress at the top of SLM specimens shows a trend that decreases with the reduction in height, which then converts into compressive stress (Vierneusel et al. 2017; Mao et al. 2017). The specimens were cut at 1 mm below the top surface to avoid the influence of residual tensile stress. To evaluate the laser energy input on the different melting modes, the E was used (Ghayoor et al. 2019):(1) $E=P/vht$(1) E was 42.03 J/mm³ in CM and 83.33 J/mm³ in KM.

The relative density of the specimens was determined by employing the Archimedes’ method. The phases of the powder and SLM specimens were analysed using a D/Max-2250 X-ray diffractometer (XRD) with Cu-Kα radiation. Measurements were performed by step scanning 2θ from 10° to 90° with a scanning rate of 5°/min. Metallographic specimens were prepared following standard procedures and etched with HNO₃: HCl: HF: H₂O (5: 3: 2: 190) solution for 20 s. The microstructures of SLM specimens were examined using an optical microscope (OM, Leica DM2700P), a field emission gun scanning electron microscope (SEM, Nova NanoSEM 230) equipped with an electron back-scatter diffraction (EBSD) detector, and a transmission electron microscope (TEM, FEI Talos F200X). X-ray energy dispersive spectroscopy (EDS) and high angle annular dark-field scanning mode was employed in STEM mode (HAADF-STEM) to characterise the cellular-like structure of the KM specimen in detail. The surfaces of the specimens after wear tests were examined using a laser confocal scanning microscope (LCSM, OLYMPUS OLS5000). Nanoindentation measurements were conducted at a load of 30 mN with a holding time of 15 s. The dynamic nanohardness, H, was determined from (Oliver and Pharr 1992):(2) $H=\frac{P_{max}}{A_{\mathrm{c}}}=\frac{P_{max}}{26.43h_c^2}$(2) where P_max is the peak load and h_c is the contact depth under the peak load. The reduced Young’s modulus (E_r) of the specimen is extracted from the effective modulus as follow (Medeiros et al. 2015):(3) $\frac{1}{E_r}=\frac{1-v_s^2}{E_s}+\frac{1-v_i^2}{E_i}$(3) where E_s and v_s are Young’s modulus and Poisson’s ratio of the specimen (v_s refers to 0.33 (Sol et al. 2018)), and E_i and v_i are Young’s modulus and Poisson’s ratio of the indenter (1140 GPa and 0.07, respectively). Dry sliding wear tests were performed using a HT-1000 ball-on-disk type tribometer (ZhongKe KaiHua, Inc., China) at room temperature. The wear disks were held against a GCr15 bearing steel ball (radius = 2 mm, HRC 60) with a load of 10 N, while rotating at a constant speed of 560 r/min for 20 min (rotation radius = 2 mm). The volumetric wear loss (V_loss) can be calculated from V_loss  = M_loss/ρ, where M_loss is the mass loss of the specimen, and ρ is the theoretical density of the AlSi10Mg (2.68 g/cm³). Thus, the wear rates (ω) can be calculated by ω = V_loss/(NS), where N is the applied load and S is the total sliding distance.

3. Results and discussion

3.1. Powder analysis

Figure 1(a) shows typical secondary electron (SE) images of the AlSi10Mg powder. The powders are spherical shape with only a few satellites. Gas pores are only occasionally observed inside the powder (inset of Figure 1(a)). The powder particle size distribution was determined using a laser particle size analyser (LPSA): Figure 1(b) shows a typical log-normal size distribution with particle sizes of 9.3 μm (D₁₀), 25.3 μm (D₅₀) and 51.0 μm (D₉₀), where D_X refers to the cumulative size distribution up to X% (including this value) of the total volume of powder.

Figure 1. (a) SE images of AlSi10Mg powder (inset shows cross-sectioned powders), and (b) particle size distribution.

3.2. Surface morphologies and densification

Figure 2 shows typical surface morphologies (top view) of SLM specimens fabricated using the two different melting modes. The surface of the CM specimen was somewhat rough with a balling phenomenon (Figure 2(a)). This morphology is due to the high viscosity and low wettability of the melt induced by the low E (Guo et al. 2019). When E is increased to that used for KM (Figure 2(b)), the surface becomes smoother along with the appearance of an irregular border. This indicates vaporisation of Al occurred. Although the accuracy of part dimensions emerges as a new issue, the surface quality of KM is better than that of CM. The relative density of SLM part under CM (96.68 ± 0.06%) is slightly higher than that of the part under KM (95.52 ± 0.05%), implying that the laser-induced balling effect is not the only reason affecting the density of SLM parts. It has been shown that there are two major defects that have been demonstrated to be the main reason for the density reduction (Wu et al. 2017). One is the high concentration of hydrogen pores caused by the humid environment and the high E; the other is keyhole defect induced by high-energy vaporised metals.

Figure 2. OM images showing typical surface morphologies of SLM specimens fabricated in two melting modes: (a) CM, and (b) KM.

3.3. Microstructure

Figure 3 presents the top and side view OM images of the SLM AlSi10Mg for the two melting modes. The top view of the CM specimen (Figure 3a) exhibits semi-cylindrically-shaped contours with an angle of 67°, which is consistent with the scanning strategy. The corresponding side view (Figure 3c) demonstrates well-overlapped molten pools characterised by a semi-circular shape owing to the applied h (120 μm) and t (50 μm). The overlapping morphology of the molten pool indicates the complete fusion of the powder particles and strong bonding within the layers. However, it is found that gas and hydrogen pores are present inside the molten pool.

Figure 3. OM images of SLM AlSi10Mg specimens fabricated in two melting modes: (a) and (c) CM, and (b) and (d) KM.

The top view of the KM specimen (Figure 3(b)) shows similar overlapped laser tracks while the molten pool shape in side view (Figure 3(d)) changes from shallow to deep. Note that the golden V-shaped molten pool (referred as keyhole molten pool, KMP) induced by the keyhole effect only appears in the KM specimen. Moreover, the pores located at the bottom of the KMP (known as keyhole defects) confirms that material vaporisation occurred as a result of the keyhole effect (Wu et al. 2017). By comparing the pore defects in two melting modes, it is thought that the hydrogen pores are the main factor affecting the densification. The dark contrast from the molten pool core and bright contrast from the molten pool boundary (MPB) can be seen in both top and side view OM images, and arise from the Gaussian distribution of the laser energy.

XRD patterns of the powder and SLM specimens are presented in Figure 4. All the peaks on the X-ray diffraction patterns fit either Si or α-Al phases. The diffraction peaks corresponding to Si are barely present on the diffraction patterns from the SLM specimens, indicating that the Si has largely dissolved into a supersaturated in α-Al matrix. This is presumably due to the high cooling rate of the SLM process (Xi et al. 2020). A careful comparison of the diffraction peaks showed that the intensities of Si peaks increase slightly as E was increased from 42.03 to 83.33 J/mm³, indicating a higher Si content in the KM specimen. This may be attributed to the vaporisation of Al at the higher temperature (Kang et al. 2016). Besides, the precipitation of Si upon cooling and re-precipitation of Si by the heat conduction from the following scans are much more severe in KM specimen owing to the higher processing temperature. It is also responsible for a higher Si content in KM specimen characterised by XRD apart from the vaporisation of Al.

Figure 4. XRD patterns from the alloy powder and the SLM specimens.

Figure 5 shows the top and side view SE images of the SLM AlSi10Mg specimens fabricated in the two melting modes. The inset bright field TEM images show the microstructures at higher magnification. The size of cellular-like networks in molten pool is apparently different from that in the MPB. Figure 5(a,c) reveal that the MPB of the CM specimen consists of the remelted zone (RZ) and the heat-affected zone (HAZ). The size of cells in the RZ is coarser than that in the molten pool. The cells are almost replaced by granular microstructure in the HAZ, which indicates that the microstructure in the MPB is inhomogeneous. The cells in the top view images (inset of Figure 5(a)) are composed of primary α-Al surrounded by ∼500 nm diameter Si networks. These equiaxed cells are observed to be elongated from the side view, as shown in Figure 5(c). Furthermore, the elongated cells grow perpendicular to the direction of the MPB, since the direction possesses the highest temperature gradient. This is known as preferential growth, indicating that the temperature gradient is the dominant reason for the various microstructures (Trevisan et al. 2017).

Figure 5. SE and inset BF TEM images of SLM AlSi10Mg specimens for the two melting modes: (a) and (c) CM, and (b) and (d) KM.

In contrast, from Figure 5(b,d), the KMP exhibits an ultrafine microstructure both in the transverse and longitudinal sections owing to the keyhole effect. The absence of RZ and HAZ in the MPB of the KMP indicates a sharp transition between general molten pool (GMP) and KMP. This sharp transition is mainly attributed to the formation mechanism of the KMP (evaporation-solidification) and corresponding extremely high cooling rate (10⁸ K/s) under KMP melting (Guraya, Singamneni, and Chen 2019). Without considering the effects of defects (e.g. pore and crack), the MPB is the weakest region of the SLM part (Xiong et al. 2019). The reduction of the MPB region, therefore, is beneficial to the improvement of mechanical properties. Generally, the depth of the KMP is deep enough to pass through the GMP and reach its MPB, as clearly shown in Figure 3(d). This explains the fact that the coarse elongated cells present surrounding the KMP (Figure 5(d)). Note that no keyhole defects are observed at the bottom of the KMP. This means that the keyhole defects in the KMP resulted from material vaporisation can be eliminated if the melt reflows sufficiently. Additionally, the strong convection triggered by the melt reflow and Marangoni convection lead to the transformation of cellular-like structures to Si nanoparticles (bright phases, as marked in the inset of Figure 5(d)) at the bottom of the KMP. This suggests that the morphology of the eutectic-Si phase is affected not only by temperature gradient, but also by melt convection.

Figure 6 presents an EBSD inverse pole figure (IPF) and corresponding image quality (IQ) map from the top view of the KM specimen. This clearly shows two types of MPB as mentioned above. One is induced by laser overlapping and re-melting, which is composed of RZ and HAZ with fine-grained region; the other is formed by material evaporation-solidification, which is generally located inside the molten pool. Note that the grain growth beyond the MPB (as marked in Figure 6(a)) indicates the existence of heterogeneous nucleation and epitaxial growth. Interestingly, coarse cells are present in the MPB within the fine-grained region, while the coarse-grained region in molten pool has small cells. Since the cell size is significantly smaller than the grain size, this unique microstructure is subgrains. The corresponding IQ map (Figure 6(b)) illustrates that the KM specimen contains a large fraction of low-angle grain boundaries (LAGBs, 2–10°, ∼54% of total grain boundaries), which is responsible for the improved strength (Wang et al. 2017). Accordingly, the nonuniform distribution of LAGBs in the molten pool denotes the weakness of the MPB region.

Figure 6. (a) Inverse pole figure (IPF) map and (b) image quality (IQ) map on the top view of the KM specimen.

 Figure 7 shows an HAADF-STEM image and EDS maps of the cells in the KMP. The cells are composed of the α-Al phase decorated with ∼200 nm diameter eutectic Si networks at the boundaries. The presence of these two phases is consistent with the XRD results in Figure 4. A closer look at the Si distribution in EDS mapping reveals that some Si precipitates exist inside the cells. Unlike GMP, no distinct segregation of Mg and Fe elements is observed in KMP owing to the very high cooling rate (Hadadzadeh, Amirkhiz, and Mohammadi 2019). It is proposed that the strong convection of the melt can effectively reduce the segregation of Mg and Fe in cell boundaries.

Figure 7. HAADF-STEM image of the cells in KMP and corresponding EDS mapping of the main elements (Al, Si, Mg and Fe).

 Figure 8 shows the schematic representation of melt flows, hydrogen pores, cooling rate, and keyhole dynamics during SLM processing. The temperature in the centre of the molten pool is higher than that in the boundary. Normally, the density of the melt decreases with increasing temperature. The density difference of the melt leads to buoyancy convection under the action of gravity (Riley and Neitzel 1998). In addition, in the absence of surfactant, surface tension γ has the same trend as the density changes, that is, ∂γ/∂T < 0. Thus, the shear stress, which from the centre of the molten pool to its boundary, generated by surface tension gradient also leads to Marangoni flow (Tsotridis, Rother, and Hondros 1989), as marked by black arrows in Figure 8(a). The large temperature gradient during SLM processing induces a high cooling rate (region 1), and leads to Marangoni flow transforming into strong Bénard-Marangoni convection (Mizev and Schwabe 2009). These give rise to the formation of Si network in CM. The lower cooling rate in MPB (region 2) effectively alleviates the influence of melt convection on Si morphology. In KM, the E is higher enough to vaporise the metal and form a V-shaped keyhole under the balance between recoil pressure and surface tension. The plasma formed by laser irradiation resides inside the keyhole and moves outwards at high speed along the surface of molten pool, which intensifies the convection of the melt (Qi et al. 2017). However, this dynamic balance is broken after the laser moves away. Then the melt will collapse and backfill into the keyhole, as marked by white arrows in Figure 8(b). This opposite movement direction of the melt could result in turbulence with an ultra-high cooling rate (region 3). That’s the main reason for the deposition of nano-Si particles at the bottom of KMP. It’s worth noting that the keyhole defects will form only if the melt backfill is insufficient. In comparison, the hydrogen pores formed by the reaction between H₂O and aluminum are the main part of the defects. By drying the raw powder at high-temperature, the accompanying negative effects can be effectively avoided.

Figure 8. Schematic representation of thermal history during (a) CM and (b) KM.

3.4. Tribological behaviors

The friction coefficients versus sliding time and wear rates of both the SLM and as-cast specimens are shown in Figure 9. The measured friction coefficients are somewhat variable during the sliding process, and decreases from higher initial values (run-in period) to lower values (steady state) in all specimens (Figure 9(a)). The mean friction coefficients (μ) of the three specimens in steady state wear (6–20 min) are calculated and summarised in Table 1 (the statistical data analysis is shown in Table S1). The friction coefficient of the as-cast specimen showed large fluctuations with a mean value of 0.35 and large wear rate (10.05 × 10⁻⁴ mm³N⁻¹m⁻¹). The CM specimen shows a similar friction coefficient (0.36) but much lower wear rate (6.96 × 10⁻⁴ mm³ N⁻¹ m⁻¹). As the E reached KM values, the measured friction coefficient decreases slightly (0.32) and the wear rate decreases substantially (3.82 × 10⁻⁴ mm³ N⁻¹ m⁻¹).

Figure 9. (a) Friction coefficients versus sliding time and (b) wear rates for the SLM and as-cast specimens.

Table 1. Friction coefficients (μ) of the steady state wear.

Specimens	CM (42.03 J/mm^3)	KM (83.33 J/mm^3)	As-cast
μ	0.36 ± 0.01	0.32 ± 0.01	0.35 ± 0.03

Figure 10 shows the typical morphologies and 3D surface profiles of the worn surfaces of the SLM and as-cast specimens, as well as of the corresponding wear debris. In the as-cast specimen (Figure 10(c)), severe plastic deformation occurs on the worn surface instead of abrasive grooves, implying an adhesion-dominated wear mechanism. At the beginning of the wear test, the GCr15 ball contacts with the specimens through a quite small surface area producing a large pressure. This large local pressure at the points of asperity contacts forges metallic junctions and engagements between the surfaces. In this situation, cold welding occurs between the friction pair. At larger sliding distances, the needle-like and plate-like eutectic Si structure leads to the rise of internal stress, providing an easy path for fracture of metallic junctions (Basavakumar, Mukunda, and Chakraborty 2009). The spalling of debris reduced the friction coefficient, while the exposed rough new surface was more likely due to cold welding, resulting in the increase of the friction coefficient. This repeated welding-fracturing process is the reason for the high wear rate and high friction coefficient fluctuation of the as-cast specimens (Joshi, Singh, and Chaudhary 2019). The V_loss, from the 3D surface profile, in this region is 0.108 mm³, which is the largest among three tested materials.

Figure 10. Topography and 3D surface profiles of the worn surfaces, and SE images of debris for (a) CM, (b) KM and (c) as-cast specimens.

The worn surfaces of both the CM and the KM specimens exhibit the typical morphologies that arise from abrasive wear, viz. distinct and long grooves generated by plowing (Figure 10(a,b)). The corresponding 3D surface profiles show that the V_loss of the SLM specimens is ∼34% lower than that of as-cast specimen, indicating that the wear of SLM specimens is reduced. The nano-scale cellular Si network plays a critical role in the wear behaviour, which could alleviate the adhesion of metal and inhibit the formation of a delamination layer (Prashanth et al. 2014). Compared to the as-cast specimens, this unique ultrafine microstructure also increases the surface hardness of SLM specimens from ∼110 to ∼150 HV (Girelli et al. 2017; Wei et al. 2017). The degree of repeated welding-fracture process was obviously reduced, and no adhesion was found in the SLM debris. Therefore, the friction coefficients of SLM specimens are stable, and the wear rates are lower than those of as-cast specimens. Of particular interest is that the V_loss of the KM specimen is lower than that of the CM specimen. The particles on the surface of the wear debris are considered as Si-rich oxide particles, details in Figure S1. These particles indicate that the wear mechanism is the dominant three-body wear, which is further verified by the wear coefficient K (0.5–5 × 10⁻³ for the three-body wear) calculated by Archard wear equation (Greer, Rutherford, and Hutchings 2002; Wu et al. 2013). Also, there are more Si-rich oxide nanoparticles on the surface of KM debris compared with CM debris. This is probably due to the finer Si networks (∼200 nm) and Si nanoparticles produced by the keyhole effect. The Si nanoparticles embedded in the KMP fall off the worn surface, which reduces the contact area of the friction pair during sliding, thereby improving the wear performance.

It is well accepted that hardness plays a significant role in determining the wear resistance of materials (Deuis, Subramanian, and Yellup 1997; Yusoff and Jamaludin 2011). In fact, the wear of materials is related to the H/E_r ratio (Pintaude 2013). A high H/E_r ratio normally denotes good wear resistance. Further, the ratio $H^3/E_r^2$ is a critical parameter to reveal the material’s ability to resist plastic deformation under a load (Herrera-Jimenez et al. 2019). A higher $H^3/E_r^2$ ratio normally indicates that the material has better resistance to plastic deformation. Hence, these parameters are suitable for evaluating the anti-wear ability of materials, as the gradual removal of material during the wear process is associated with plastic deformation (Ehtemam-Haghighi, Cao, and Zhang 2017). The H and E_r of two molten pools are measured using nanoindentation tests, as shown in Figure 11. Figure 11(b) depicts the nanoindentation load-depth curves collected from two types of molten pools (marked in Figure 11(a)). The curves are smooth without observable serration. After unloading, there is obvious permanent deformation and elastic recovery. The corresponding average values of H and E_r are presented in Figure 11(c). It is found that KMP induced by evaporation shows higher H and lower E_r than those of the GMP. Further, the H/Er and $H^3/E_r^2$ ratios of the two molten pools exhibit a similar trend (Figure 10(d)), suggesting that the existence of the KMP is mainly responsible for the improved anti-wear ability.

Figure 11. (a) Nanoindentation positions in the two molten pools from the top view in the KM specimen; (b) corresponding load-displacement curves; (c) the values of H and E_r; and (d) H/E_r and $H^3/E_r^2$ ratios.

4. Conclusions

The following conclusions can be drawn from this experimental study:

1. Although both hydrogen pores and keyhole defects can reduce the density of SLM AlSi10Mg, the former is the dominant factor. The keyhole defects in keyhole molten pool (KMP) can be eliminated if sufficient melt reflow occurs.

2. The increase of the volumetric energy density (E) does not change the nominal composition of SLM AlSi10Mg, but modifies the distribution of Si via materials vaporisation and melt reflow. Compared with the general molten pool (GMP), the molten pool boundary (MPB) in the KMP presents a sharp transition without the remelted zone (RZ) and the heat-affected zone (HAZ). The cell size in the KMP is much finer (∼200 nm) than that in the GMP (∼500 nm).

3. The wear mechanisms of the as-cast specimen are dominated by adhesive wear, while the main mechanism for the SLM specimen is abrasive wear due to the Si network.

4. The randomly distributed Si nanoparticles at the KMP reduces the contact area of the friction pair, resulting in higher anti-wear ability. The nanoindentation measurement results show that the KMP has higher H/E_r and $H^3/E_r^2$ ratios, which also prove that the existence of KMP is the main reason for better wear performances of KM specimens.

Acknowledgements

The authors would thank Sinoma Institute of Materials Research (Guang Zhou) Co., Ltd. for the assistance in TEM characterisation.

Disclosure statement

No potential conflict of interest was reported by the author(s).

Additional information

Funding

The research was supported by the National Key Research and Development Program of China [grant number 2016YFB1100103], National Natural Science Foundation of China [grant number 51771233], China Postdoctoral Science Foundation [grant number 2018M633164] and National Science Fund for Distinguished Young Scholars [grant number 51625404].

Notes on contributors

Hong Wu

Hong Wu is a professor of Central South University.

Yaojia Ren

Yaojia Ren is a currently PhD candidate of Central South University.

Junye Ren

Junye Ren is a master's degree holder in material science and engineering with Central South University.

Anhui Cai

Anhui Cai is a professor of Hunan Institute of Science and Technology.

Min Song

Min Song is a professor of Central South University.

Yong Liu

Yong Liu is a professor of Central South University.

Xiaolan Wu

Xiaolan Wu is a professor of Beijing University of Technology.

Qingxiang Li

Qingxiang Li is in the Shenzhen Zhong Jin Ling Nan Nonfemet Co., Ltd, China.

Weidong Huang

Weidong Huang is a professor of Northwestern Polytechnical University.

Xiaoteng Wang

Xiaoteng Wang is in the Lancaster University.

Ian Baker

Ian Baker is a professor of Dartmouth College.