Effect of Trace Addition of Ceramic on Microstructure Development and Mechanical Properties of Selective Laser Melted AlSi10Mg Alloy

Spherical gas-atomized AlSi10Mg powder (99.8% purity) ranging from 20 μm to 63 μm and TiB₂ powder (99.9% purity) with a size of 3‒5 μm (Figure 1a) were served as raw materials. The composite powder components consisting of 1.0 wt.% TiB₂, which were prepared by high energy ball milling in a Pulverisette 6 planetary mono-mill (Fritsch GmbH, Germany) at a rotation speed of 200 r/min for 4 h. After ball milling, TiB₂ particles were homogenously dispersed around the surfaces of the AlSi10Mg particles, as shown in Figure 1b. At the same time, it can be seen that the ceramic particles have been fragmented to a certain extent.

The as-used device in this study is SLM-150 developed by Nanjing University of Aeronautics and Astronautics (NUAA). The SLM system consists of an IPG Photonics Yitterbium YLR-500-WC fiber laser with a maximal laser power of 500 W and a laser spot diameter of 70 μm, an inert Ar gas protection system, an automatic powder layering system and a computer control system. Initially, an aluminum substrate was set on the building platform and leveled. The building chamber was sealed and Ar gas was pumped into it. Then, the TiB₂/AlSi10Mg composite powders were uniformly spread on the substrate using the automatic powder spreading system until a layer thickness of 50 μm was obtained. The process was repeated until the cubic samples were constructed with a dimension of 10 mm × 10 mm × 8 mm (Figure 2a). An “island scanning strategy” was applied to fabricate specimens, and the orientation of the laser was changed by 37° between successive layers (Figure 2b). The adoption of the island deposition strategy aims to balance the residual stresses in the build [29]. Based on previous experiments, a series of SLM experiments using a wide range of processing parameters, i.e., laser powers of 300‒400 W and scan speeds of 2000‒3200 mm/s, were performed to testify for the SLM process ability.

The densities of the SLM-processed specimens were measured using the Archimedes principle. Metallographic samples were prepared as per the standard procedures and then etched with Keller’s reagent (HF (1.0 mL), HCl (1.5 mL), HNO₃ (2.5 mL), and distilled water (95 mL)) for 30 s. The surface morphology and microstructure were characterized using an optical microscope (OM) and an S-4800 field emission scanning electron microscope (FE-SEM, Hitachi, Tokyo, Japan) at a voltage of 5.0 kV. The chemical compositions were determined by an EDAX energy dispersive X-ray spectroscope (EDX) (EDAX, Inc., Mahwah, NJ), using a super-ultrathin window sapphire detector. The phases of the specimens were identified through a Bruker D8 Advance X-ray diffractometer (XRD) (Bruker AXS GmbH, Karlsruhe, Germany) with Cu Kα radiation at 40 kV and 40 mA, and the continuous scan mode was used with a scan rate of 4°/min.

The Vickers hardness was measured using an HXS-1000AY microhardness tester (AMETEK, Shanghai, China) at a load of 0.2 kg and a dwell time of 15 s. Tensile tests using the machine (MTS E45.105) with 0.2 mm/s test speed were performed in accordance with GB/T 228.1-2010 standard at room temperature. Ultimate tensile strength, yield strength and elongation of TiB₂/AlSi10Mg composite and AlSi10Mg alloy samples were achieved. Fractographs of specimens were analyzed by SEM.

Figure 3 shows the variation in the densification behavior of the SLM produced TiB₂/AlSi10Mg composites according to different processing parameters. It is obvious that higher relative densities were obtained under laser power (P) of 400 W. With different laser scan speeds (v), the relative densities also exhibited distinct variations. At a high v of 3200 mm/s, a sample with a relative density of 95.9% was obtained. As the v decreased to 2800 mm/s and 2400 mm/s, relative densities were 97.4% and 98.6%, respectively. Near fully dense part with a relative density of 99.4% was obtained at the lowest v of 2000 mm/s. Thus the relative densities were increased as the v decreased under the same P. As the P decreased, the relative densities of the samples decreased accordingly. When the P decreased to 300 W, the relative densities decreased sharply as the v increased, implying a high sensitivity to scan speed.

The typical surface morphologies of the SLM-processed TiB₂/AlSi10Mg parts using various laser scan speeds at a laser power of 400 W are displayed in Figure 4. At a relatively high v of 3200 mm/s, an obvious uneven surface containing a large number of irregular-shaped balls and pores was generated. As the v decreased, the balling effect was still present but the pores were reduced clearly, leading to the surface quality improved. When the v reached the minimum value of 2000 mm/s, a relatively smooth surface was obtained with the highest relative density.

The densification behavior is intrinsically associated with the viscosity of high temperature metal melt, wettability and liquid-solid fluid characteristics [30]. In the SLM molten pools, the relationship between dynamic viscosity μ and temperature T can be expressed as:where m is the atomic mass, k is the Boltzmann constant, and γ is the surface tension of the melt. A low scan speed means a long dwelling time of laser beam and high energy input, and the resultant high T results in a decrease of μ. The low viscosity of the liquid phase can improve the liquidity of the metal melt, which also contributes to the wettability between TiB₂ reinforcements and the Al matrix.

Generally, the balling effect is a typical metallurgical defect in the SLM process, which closely related to the temperature of the liquid phase, wetting ability and Marangoni convection [31, 32]. At a high scan speed, the laser tracks easily show a shrinking trend for reducing the surface tension due to the poor wetting characteristic. Through Arafune and Hirata’s results [33], the intensity of Marangoni convection can be estimated by dimensionless Marangoni number (Ma) as follows:where Δσ is the surface tension difference of Marangoni flow, L is the length of the free surface, μ is the dynamic viscosity, and v_k is the kinematic viscosity. It can be concluded that the convection can be reinforced by the decrease of μ within the pools. With the decrease of scan speed, Marangoni convection is intensified as the viscosity decreases according to Eq. (2), which improves the spreading ability of the molten pool. Therefore molten tracks present a better continuity without balling phenomenon, owing to the well-wetting characteristic [34].

The different interlayer microstructures with different scan speeds are visible upon etching on cross-sections of SLM-processed TiB₂/AlSi10Mg parts (Figure 5). At a high scan speed of 3200 mm/s (Figure 5d), a significant poor interlayer bonding can be observed including several considerable inter-layer pores, which was in line with the density results. As the scan speed decreased to 2800 mm/s (Figure 5c), some small pores still existed, and cracks can be identified in the molten pool. As the scan speed further decreased to 2400 mm/s (Figure 5b), there were no obvious pores and the cross-section showed a relatively heterogeneous layerwise microstructure, indicating the continuity of scanning trace was poor. When the lowest scan speed of 2000 mm/s was used, the molten pool showed a more coherent inter-layer bonding free of residual pores (Figure 5a). The molten pools had a clear configuration, showing a stable solidification process during SLM. The different solidification behaviors are believed to be controlled by the laser energy input and fluid flow conditions. Higher viscosity and surface tension could block the flowability of the melt and reduce the overall rheological performance of the composite melt, thus make it difficult for the liquid phase to expand sufficiently. Simultaneously, balling hinders the spread of the new layer powder on the previous layer, resulting in the formation of inter-layer pores.

Typical XRD characterization corresponding to the α-Al phase of SLM-processed TiB₂/AlSi10Mg parts at different processing parameters is depicted in Figure 6. It is obviously that the 2θ locations of the diffraction peaks for α-Al significantly changed as the variation of applied SLM parameters. Generally, the standard diffraction peaks for the α-Al phase are located at 2θ = 38.4° and 2θ = 44.7° (JCPDS Card No. 04-0787). The corresponding quantitative parameters for composites obtained from XRD analysis are listed in Table 1. The specific 2θ locations, intensities and FWHM (Full Width at Half Maximum) of α-Al peaks show obvious changes on account of the addition of TiB₂ and the variation of SLM processing parameters. With the decrease of scan speed, the location of α-Al diffraction peaks shifted to higher 2θ angles. Based on Bragg’s law for XRD analysis [35]:

the observed increase of 2θ revealed that a decrease in lattice plane distance, implying the crystal lattices were distorted, which was responsible for a small amount of Si and partial micro-TiB₂ particles were dissolved and supersaturated in the Al matrix. The atomic radius of Al, Si, Ti and B are 0.1431 nm, 0.1172 nm, 0.1453 nm and 0.0950 nm, respectively. The replacement solid solution usually generates in α-Al matrix [36], resulting in lattice distortion. With the scan speed decreases, the laser-induced non-equilibrium metallurgical process is reinforced which could dissolve more Si atoms [37]. It is clearly that the atomic radius of Si is smaller than Al, causing the decrease of the lattice plane distance. Table 1 showed the value of FWHM decreased as the v increased. According to Scherrer’s formula [38], the increase of FWHM reveals the crystalline size reduction of the composite. The high scan speed leads to the rapid cooling rate, thus the grain growth is prevented, which is sensitive to the SLM processing parameters.

The SEM image and corresponding EDX mapping analysis reveal the distribution state of the TiB₂ reinforcing particles at v=2000 mm/s that obtained the highest relative density (Figure 7). Figure 7a shows the homogeneous distribution state of TiB₂ particles. To further elucidate the particles were TiB₂, the corresponding element EDX mappings of the Al, Si, Mg, Ti and B are given in Figure 7b‒f, respectively. Obviously, these particles were TiB₂ phase as represented by the concentration of Ti.

In the solidification microstructures generated by traditional casting [39], the majority of TiB₂ particles aggregate easily. However, in this study, the clustered particles were dispersed effectively, which could be contributed to the ball milling procedure and SLM process. For the low scan speed, the enhanced Marangoni convection and flowability of the melt promote the movement of TiB₂ particles, improving their uniform distribution and the wettability between reinforced particles and the matrix. Furthermore, particle pushing and engulfment can not be ignored in particulate-reinforced metal matrix composite systems, which contributes to a relatively uniform distribution of particles [39].

The morphology of TiB₂ particles in SLM-processed composite is illustrated in Figure 8a. Compared to the initial powder (Figure 1a), it could be observed that the TiB₂ particles became more round and smooth after the SLM process, indicating that they are partially melted. No obvious reaction products were observed, implying the melted TiB₂ dissolved into the matrix. This is consistent with the XRD results. Generally, the laser absorption efficiency of ceramics is better than that of metals [18]. Sufficient laser energy input makes the heat concentrate on the TiB₂ easily, causing the melt of TiB₂ particles although they have a high melting point (3225 °C [40]). The sharp edges of the particles tend to concentrate energy, where melt preferentially and generate smooth edges.

It has been reported that TiB₂ particles and α-Al grains tend to form a high-coherency orientation relationship between the two atomic structures to reduce the solid-particle interfacial energy [39]. The smooth TiB₂ particles in uniform dispersion and attendant enhanced wettability between reinforcement and aluminum matrix contribute to the formation of coherent interfaces. There were no palpable voids or reaction products can be seen near the interface, indicating that the interface between the TiB₂ particles and the matrix was well integrated. Further, a small amount of spherical particles could be seen precipitating at the interface (see the rectangle in Figure 8a). Figure 8b shows a high-magnification SEM image from the local region of Figure 8a, where coarse particles at the interface are shown more clearly. For the micro-scale TiB₂, there has not been evidentially reported reacting with the aluminum matrix. However, eutectic Si tends to form a rod- or needle-like coarse microstructure (usually tens to hundreds of microns in size) in conventional casting Al alloy [41]. It is well known that the rapid cooling process can lead to significant extended (non-equilibrium) solid solubility in alloys [42, 43]. Thus the SLM process can dissolve more Si phase compared with the traditional casting technology due to its extremely high cooling rate. With the low scan speed of 2000 mm/s, the cooling rate decreased and the eutectic Si phase has ample time to precipitate from the Al matrix. The TiB₂ particles affect the distribution of Si particles in the matrix. On the one hand, it has been reported that boride particles with spherical morphology would increase the chance of exposing crystallographic plane suitable for precipitation [44]. On the other hand, the presence of TiB₂ particles could cause chemical inhomogeneity, and these particles can act as good heterogeneous nucleation sites for the formation of Si particles. Therefore, it can be concluded that these nano-particles precipitated at the interface are probably Si particles.

Figure 9 illustrates the microhardness and its distribution measured along the building direction of SLM-processed AlSi10Mg alloy and TiB₂/AlSi10Mg composite parts. Compared with the microhardness of AlSi10Mg alloy, the average microhardness of the composite (120.0 HV) was slightly higher than that of the alloy (117.5 HV). Certainly, compared with cast AlSi10Mg (95 ~ 105 HV) [45], the microhardness have been both greatly improved. This is attributed to the large distribution of the Si phase in the Al matrix due to the rapid solidification and cooling rate of SLM. Furthermore, the microhardness of the AlSi10Mg part showed a relatively large fluctuation varied from 110.2 HV to 124.4 HV. The values also increased as it was closer to the bottom of the sample. The closer it is to the bottom of the sample, the greater the cooling rate and temperature gradient of the melt will be, causing the larger residual stress, which increases the microhardness. The heat accumulation near the top surface could result in stress release, leading to the reduction of microhardness [45]. The composite part showed a reduced fluctuation of the microhardness values. This can be attributed that the uniformly dispersed TiB_2. particles relieve the stress release and effectively restrain the deformation of the matrix during indentation, hence stabilize and elevate the microhardness value.

The mechanical properties of the TiB₂/AlSi10Mg composite and AlSi10Mg alloy fabricated by SLM were investigated by room temperature tensile tests and the results are provided in Figure 10. The ultimate tensile strength (UTS), yield strength (YS) and elongation of the composite were 380.0 ± 10 MPa, 250.4 ± 6 MPa and 3.43 ± 0.2%, respectively, increased by 9.73%, 64.4% and 4.57% compared to the SLM-processed AlSi10Mg alloy (346.3 ± 8 MPa, 152.3 ± 4 MPa and 3.28 ± 0.15%) under the same processing parameters. It can be seen that both strength and toughness were improved. Figure 11 illustrates the morphologies of the fracture surfaces for the AlSi10Mg part and the TiB₂/AlSi10Mg composite part. Some spherical pores on the fracture surface of the AlSi0Mg could be observed clearly (Figure 11a), these porosities maybe the origin of crack initiation and lead to the poor mechanical properties [46]. The fracture surface of the composite shows a flat fracture surface free of macro-pores (Figure 11c). Moreover, the fracture surfaces at high-magnification were both apparently covered with dimples, indicating their plastic fracture characteristics (Figure 11b and d). Obviously, the dimples of TiB₂/AlSi10Mg composite were finer and more uniform than AlSi10Mg alloy, revealing its enhanced plasticity and excellent ability to absorb the energy of the fracture.

It is known that the distribution of the reinforcing particles in the matrix and the degree of bonding between the matrix and reinforcement are the key factors for composites’ mechanical properties [41]. It has been proved that TiB₂ particles dispersed uniformly and a strong interfacial bonding was generated (Sections 3.3 and 3.4), which is beneficial for the coordinate deformation between the strengthening particles and the matrix. The smooth edge without a sharp angle of TiB₂ could release the stress concentration and restrain the cracks under external loads. Based on our analysis above, the increment in tensile properties of SLM processed TiB₂/AlSi10Mg composite is attributed to these three reasons: (I) The dissolved Si and TiB₂ phases result in lattice distortion of Al, forming local stress fields which could interact with dislocations and restrain their motion. Hence, the tensile strength is improved. (II) The uniformly dispersed micro-TiB₂ particles effectively inhibit the slipping of dislocation during deformation via the Orowan mechanism. (III) The strong interfacial bonding between the micro-particles and the matrix facilitates the transfer of the load from the aluminum matrix to TiB₂ particles [41, 45, 47‐49]. The strengthening effect caused by load transfer on yield strength can be specifically expressed by the shear lag model modified by Nardone and Prewo [50]:where σ_c and σ_m are the yield strengths of composite and the alloy (~152.3 MPa), respectively. V_p and V_m are the volume fraction of the reinforcing phase and matrix (~1 vol% and ~99 vol%), s is the aspect ratio of TiB₂ (~2.5). As such, the calculate σ_c is 153.3 MPa. It contributes more to the improvement of yield strength. Orowan strengthening via dispersed TiB₂ particles in the matrix is expressed by the equation [50, 51]:where G is the shear modulus of the matrix (2.59 × 10⁴ MPa), b is the Burger’s vector (0.286 nm), d_p denotes the average particle diameter (4 μm). The estimated strength from the Orowan strengthening is just 0.246 MPa. It is worth noting that this contribution is negligible compared with load transfer strengthening. Solid solution strengthening has been proved to exist in both alloy and composites [27, 45]. In this study, the solid solution strengthening is weakened due to the precipitation of the Si phase, and it is inappreciable compared with alloy for composite. Therefore, load-bearing is the major strengthening mechanism for micro-TiB₂/AlSi10Mg composite.