Anomalous strengthening by supersaturated solid solutions of selectively laser melted Al–Si-based alloys

Additive Manufacturing 33 (2020) 101152

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Anomalous strengthening by supersaturated solid solutions of selectively laser melted Al–Si-based alloys

T

Naoki Takata*, Mulin Liu, Hirohisa Kodaira, Asuka Suzuki, Makoto Kobashi Department of Materials Process Engineering, Graduate School of Engineering, Nagoya University, Furo-cho, Chikusa-ku, Nagoya, 464-8603, Japan

A R T I C LE I N FO

A B S T R A C T

Keywords: Additive manufacturing Aluminum alloy Strain hardening Rapid solidification Transmission electron microscopy

To identify the dominant contributing factor in the anomalously high strength of Al–Si-based alloys fabricated by selective laser melting (SLM), microstructural characteristics of a SLM-built Al–10Si–0.3 Mg alloy (AlSi10Mg) and their changes upon annealing at elevated temperatures were investigated. The as-built AlSi10Mg alloy exhibits a peculiar microstructure comprising of a number of columnar α-Al (fcc) phase with concentrated Si in solution. Numerous nano-sized particles were observed within the α-Al matrix. At elevated temperatures, a number of Si phase (diamond structure) precipitates consumed the solute Si in the columnar α-Al phase, but the microstructure of the α-Al matrix changed slightly. After annealing at elevated temperatures, the tensile strength of the as-built AlSi10Mg alloy substantially decreased accompanied by a reduction in the strain hardening rate. The supersaturated solid solution of the α-Al phase containing numerous nano-sized particles enhanced the strain hardening, resulting in the anomalous strengthening of the SLM-built AlSi10Mg alloy. The microstructural features were formed due to rapid solidification at an extremely high cooling rate in the SLM process, which provides important insights into controlling the strength of Al–Si-based alloys fabricated by SLM.

1. Introduction Additive manufacturing (AM) of metals represents a promising route to fabricate highly complex metal parts without the need for special facilities. Although metal AM technologies include various processing routes [1,2], one of the most common routes used for metals and alloys is powder bed fusion (PBF) [3,4]. PBF involves the use of either laser or electron beams to melt and fuse powder particles of metals or alloys. This process includes the commonly used selective laser melting (SLM), selective laser sintering (SLS), direct metal laser sintering (DMLS), and electron beam melting (EBM) [1–4]. Recent developments in the PBF technologies are being applied to the SLM process [3,4] to enable the manufacturing of complex geometrical parts using various metals and alloys including lightweight aluminum (Al) alloys. It is generally known the SLM process produces bulk Al alloys with superior strength to those of conventionally produced alloys. Fig. 1 shows the strength levels of the SLM-built Al–Si-based alloys [5–13] (on the vertical axis) were plotted as a function of the strength of those produced by conventional gravity or die-cast cast processes [13–15] (on the horizontal axis). The SLM-fabricated alloys exhibit higher strength than the conventionally cast alloys (Fig. 1(a)). It is intriguing that the difference in strength between SLM-fabricated and conventionally cast ⁎

alloys appears more significant in higher-strength alloys, indicating the possibility that higher-strength Al–Si-based alloys can be made much stronger by SLM processing. This anomalous strengthening could be due to the characteristic microstructures in locally melted and rapidly solidified alloy samples produced via the SLM process [16,17]. It is noteworthy that conventional heat treatments (T6: solution treatment and subsequently artificial aging) often reduce the strength of SLMfabricated alloys, resulting in the same strength level to heat-treated conventionally cast alloys (Fig. 1(b)). These data show that the SLMfabricated alloys experience a loss in strength on exposure to conventional heat treatments. It is therefore essential to identify any microstructural characteristics (providing high strength) produced by the SLM process. In order to understand the mechanism involved, significant strain hardening can be considered as a remarkable mechanical behavior of SLM-fabricated Al–Si-based alloys [7,6–13]. A number of literature studies have reported significant strain hardening that results in the high strengthening of SLM-fabricated alloys. The mechanisms of the unique strengthening were proposed as an Orowan looping mechanism involving fine particles of eutectic Si phase [7], supersaturated solid solutions with solute impurity elements [11], Mg2Si precipitation hardening [18], and multiple hardening mechanisms [19] in the SLMfabricated Al–Si-based alloys, whereas the dominant contributing factor to the significant strain hardening still remains unclear. Wu et al. [20]

Corresponding author. E-mail address: [email protected] (N. Takata).

https://doi.org/10.1016/j.addma.2020.101152 Received 10 September 2019; Received in revised form 1 January 2020; Accepted 23 February 2020 Available online 25 February 2020 2214-8604/ © 2020 Elsevier B.V. All rights reserved.

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Fig. 1. Tensile strength of various Al–Si based alloys fabricated by selective laser melting in comparison with conventionally casted alloys: (a) as-fabricated samples (as-SLM built or as-cast), (b) subsequently heat-treated samples (T6 treatment). Table 1 Chemical compositions of the studied AlSi10Mg alloy powder and fabricated samples (wt%).

Nominal ICP analyzed

Powder As-fabricated

Si

Fe

Cu

Mn

Mg

Ni

Zn

Pb

Sn

Ti

9.0 ∼ 11.0

≤ 0.55 0.42 0.40

≤ 0.05

≤ 0.45 – –

0.20 ∼ 0.45

≤ 0.05 – –

≤ 0.10 – –

≤ 0.05 – –

≤ 0.05 – –

≤ 0.15 – –

10.73 10.77

– –

0.17 0.18

systematically characterized. The mechanical behavior and thermal conductivity of the SLM-fabricated sample were investigated to compare with those of annealed samples. A dislocation substructure was observed in the tensile-deformed specimens. These results were used to discuss microstructural factors that contribute towards the high strength of the SLM-fabricated AlSi10Mg alloy.

carried out in situ transmission electron microscope (TEM) observations on dislocation multiplications in a SLM-fabricated AlSi10Mg (Al–10Si–0.3 Mg) alloy under compression. The in situ observations showed that extremely high-density dislocations exist inside the α-Al (fcc) matrix in the deformed samples, suggesting an increase in internal stress within the α-Al grains rather than around the eutectic cell boundaries decorated with fine particles of Si phase. It has been reported that the SLM-fabricated AlSi10Mg alloy exhibited not only a high strength of approximately 470 MPa but a much low thermal conductivity of 103 W·m−1 K−1 [21]. After the heat treatment at 300 °C, its thermal conductivity became as high as 173 W·m−1 K−1 [21], comparable with that of conventional Al–Sibased alloys. Its strength was significantly reduced to approximately 300 MPa. It is intriguing that the heat treatment significantly improves the thermal conductivity of SLM-fabricated AlSi10Mg alloy, regardless of a slight change in microstructural morphology observed at SEM resolution [8]. In general, the thermal conductivity measured for bulk metallic materials is sensitive to solute elements and their concentration (in metal matrix) rather than the other microstructural parameters (second phase, grain size and dislocation density) [22]. In the aspect of microstructural features in the SLM-fabricated AlSi10Mg alloy, its low thermal conductivity suggests the presence of highly concentrated Si (or impurity elements) in solution. The solute atoms in the α-Al matrix might play a role in solid-solution strengthening, which contributes to the anomalous strengthening (as presented in Fig. 1). However, there are few reports [21] on thermal conductivities of SLM-fabricated and subsequently heat-treated Al–Si-based alloys. In particular, the correlation between microstructural features and thermal conductivity (and its associated mechanical properties) still remains unclear. In the present study, to identify the dominant contributing factor in the anomalous strengthening of Al–Si-based alloys produced by the SLM process, an AlSi10Mg alloy with a nominal composition of Al–10Si–0.3 Mg (wt.%) was investigated as a most commonly used alloy in metal AM technologies. The microstructure of the SLM-fabricated AlSi10Mg alloy and its changes at elevated temperatures were

2. Experimental In the present study, a commercially used AlSi10Mg alloy powder [21] was studied, the nominal composition and the composition measured by inductively coupled plasma atomic emission spectroscopy are summarized in Table 1. The major elements (Si, Fe, and Mg) present in the fabricated samples were found to be consistent with the nominal composition. Cube samples (45 mm in length) were built at room temperature using an EOSINT M 280 AM system (EOS GmbH, Germany) equipped with a Yb-fiber laser. High-purity Ar gas was flowed into the chamber to prevent oxidation of the built samples during the SLM process. The base plate for powder bed was not heated in the present SLM process. The processing parameters were as follows: the laser power was 380 W, the laser scanning speed used was approximately 1 m·s−1, the powder layer thickness was 30 μm, and the focus size was approximately 100 μm. The laser scanning tracks were performed using the rotation angle of 67° between each powder layer. The X/Y-plane was defined as being parallel to the powder-bed plane and the Z-direction was defined as the building direction normal to the powder-bed plane. These are described in detail elsewhere in the literature [8]. The as-fabricated bulk samples were annealed at 300 or 530 °C for various periods ranging from 60 s (s) to 48 h (h), followed by quenching in water. An atmospheric furnace was used for long-term annealing (above 600 s in holding time), while a commercial chloride salt bath was used for short-term annealing (below 600 s in dipping time). Note that thermodynamic calculations [8,23] have been used to identify that in the studied alloy composition, different temperatures of 300 and 2

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Fig. 2. Optical micrographs showing (a) as-fabricated AlSi10Mg alloy sample and (b-f) subsequently annealed samples: (b) 300 °C for 60 s, (c) 300 °C for 2 h, (d) 300 °C for 6 h, (b) 530 °C for 60 s, (c) 530 °C for 6 h.

530 °C correspond to a four-phase region of α-Al + Si (diamond) + βAlFeSi + Mg2Si and a three-phase region of α-Al + Si (diamond) + βAlFeSi, respectively. Different methods were used to prepare the samples used for the observations. The prepared bulk samples were cut into a couple of pieces and then mounted in conductive Bakelite. The mounted samples were mechanically and then electro-polished with a solution of a 1:9 vol ratio of perchloric acid and ethyl alcohol at room temperature prior to being observed by optical microscopy. The samples for the scanning electron microscope (SEM) observation from the X/Y direction were ion-polished by a cross-section polisher at 6 V. These microstructures were observed using SEM operating at 30 kV. Orientation distributions were analyzed by electron backscatter diffraction (EBSD) using a step size of 2 μm. The thin samples for the TEM observations were prepared using a JEOL ion slicer at 6 V. The TEM observations and energy-dispersive X-ray spectroscopy (EDS) analyses were performed using a JEOL JEM-2100 plus operating at 200 kV. High-resolution TEM observations were conducted using a JEOL JEM-2100 F/HK. The hardness (HV) of these samples was measured using a micro-Vickers indenter at a fixed load of 9.8 N at ambient temperature. Unidirectional tensile tests

using plate specimens with a gauge length of 14 mm and a thickness of 2 mm were performed at an initial strain rate of 1.2 × 10−3/s at ambient temperature. The tensile test specimens were cut out from the center part of the fabricated cube samples. In the present study, the tensile direction was applied to a direction perpendicular to the building direction (along the X/Y direction) in order to compare the strength level of each specimen. The basic results of preliminary tensile tests using the present plate specimens of the studied AlSi10Mg alloy have been reported [8], indicating the reproducible results of mechanical properties measured by the present tensile tests using the plate specimens. The directional dependence of the tensile properties has been previously described in detail in the literature [8]. In addition, tensile tests were interrupted at a nominal strain of approximately 3 % for a couple of specimens (prepared from the as-fabricated and 300 °C/ 2 h annealed samples). TEM was used to observe dislocation substructures in the interrupted specimens. In order to evaluate the thermal conductivity (λ: W·m−1 K−1), discshaped samples with a diameter of 10 mm and a thickness of 1.5 mm were cut from the fabricated and subsequently annealed samples. The thermal diffusivity (α : m2/s) and specific heat capacity (Cp: J·g−1 K−1) 3

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Fig. 3. Orientation color maps (obtained by EBSD analyses) for (a) as-fabricated AlSi10Mg alloy sample and (b-d) subsequently annealed samples: (b) 300 °C for 60 s, (c) 300 °C for 2 h, (d) 530 °C for 6 h.

elongated grains with a mean width of approximately 10 μm. Smaller equiaxed grains with a mean size of below 5 μm appeared along the melt-pool boundaries. The elongated grains often have a < 001 > orientation along the Z direction, resulting in the development of a {001} texture. These microstructural characteristics of the α-Al matrix (Fig. 3(a)) can be observed in the samples annealed at 300 °C (Fig. 3(b, c)). After annealing at 530 °C, the microstructural morphology remained unchanged even after 6 h (Fig. 3(d)), whereas the size of the elongated α-Al grains slightly increased. These detailed microstructural factors characterized by EBSD analyses have been described in detail in the literature [8]. Fig. 4 shows secondary electron SEM images featuring microstructures of the as-built sample and the samples annealed at 300 or 530 °C. A number of columnar α-Al cells (these tend to be elongated along the building direction) with a mean width of approximately 0.5 μm surrounded by numerous fine particles of Si phase (diamond structure) were observed inside the melt pools in the as-built sample (Fig. 4(a)). The columnar α-Al grains and surrounding fine particles of Si phase correspond to the primary solidified phase and α-Al/Si eutectic (corresponding to the Si-rich liquid with a lower melting temperature) in the rapidly solidified regions, respectively [9,27,28]. In the sample annealed at 300 °C for 60 s (Fig. 4(b)), many fine particles of Si phase were observed inside the columnar α-Al cells and some particles of Si phase appear coarser in the α-Al/Si eutectic regions. High-magnification SEM images are presented in Fig. 5. The minute observation identified fine Si phase precipitated within the columnar α-Al phase at an elevated temperature of 300 °C (Fig. 5(a, b)). These results indicate that the Si phase is present as a fine precipitate within the α-Al matrix at elevated temperatures above 300 °C. The precipitate morphology has been observed in various Al–Si-based alloys annealed at relatively low temperatures (around 300 °C) [24]. During annealing at 300 °C, the precipitates of Si phase and become coarse, resulting in the relatively homogenous distribution of these Si particles in the α-Al matrix observed in the annealed samples (Figs. 4(c,d) and 5(c)). After annealing

of prepared samples were measured using a laser flash method at ambient temperature (approximately 20 °C). The experimentally measured values of α and Cp enable the calculation of the thermal conductivity (λ) of the samples according to a simple relationship expressed as λ = α ·CP ·ρ, where ρ is the density of the AlSi10Mg alloy used (2.67 g/cm3) [21]. 3. Results 3.1. Microstructure Fig. 2 shows low-magnification optical micrographs observed from the X/Y direction exhibiting the microstructures of the as-built sample and the samples annealed at 300 or 530 °C for various periods. The Z direction (building direction) corresponds to the vertical direction in all images. The as-built sample (Fig. 2(a)) exhibits a unique microstructure comprising of a number of melt pools corresponding to the locally melted and rapidly solidified regions produced by scanning laser irradiation in the SLM process. The observed microstructural characteristics of the melt pools were found to slightly change upon annealing at 300 °C (Fig. 2(b–d)), even after 6 h (Fig. 2(d)). The observed melt-pool boundaries correspond to numerous fine Si particles localized at the interface between the solid and liquid phases (locally melted by laser irradiation). These fine Si particles appear stable at an elevated temperature of 300 °C. After annealing at 530 °C for 60 s, the morphologies of melt pools were scarcely observed in the optical micrograph (Fig. 2(e)). The finely distributed Si particles became obviously coarser after 6 h (Fig. 2(f)). These observed microstructural changes correspond well with the results of previous studies on Al–Si-based alloys [14,24–26]. Fig. 3 shows the orientation color maps prepared by EBSD analyses for the as-built sample and the samples annealed at 300 or 530 °C. The orientation along the Z-direction is presented according to the color key in the unit triangle. These EBSD analyses revealed that the microstructures of the α-Al (fcc) matrix consist of a large number of 4

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Fig. 4. SEM images showing microstructures of (a) as-fabricated AlSi10Mg alloy sample and (b-f) subsequently annealed samples: (b) 300 °C for 60 s, (c) 300 °C for 2 h, (d) 300 °C for 6 h, (b) 530 °C for 60 s, (c) 530 °C for 6 h.

equilibrium) after 6 h. These results demonstrate the presence of a certain amount of solute Si in the columnar α-Al phase of the as-built sample, which is supported by the fine Si precipitate observed within the α-Al phase of the annealed samples (Fig. 5(b, c)). Fig. 7 shows bright-field TEM images that exhibit the microstructures of the as-built sample and the 300 °C/2 h annealed sample. The observed TEM image shows a number of columnar α-Al cells in the as-fabricated sample (Fig. 7(a)), which corresponds well to the microstructural morphology observed by SEM (Fig. 4(a)). Inhomogeneous contrast variations were observed locally inside the columnar α-Al cells. Diffraction spots derived from Si phase with diamond structure in the eutectic regions were detected in the corresponding selected area electron diffraction (SAED) pattern (Fig. 7(b)). Although clear diffraction spots derived from the Si phase were not observed, the observed diffraction ring (corresponding to the reflection from {111} planes) could correspond to fine eutectic Si phase. In the annealed sample, numerous Si particles with a diameter of several tens of nanometers were observed within the columnar α-Al phases (Fig. 7(c)). Uniform contrast was observed inside the α-Al matrix, which is different from that observed in the as-built sample (Fig. 7(a)). The SAED pattern (Fig. 7(d)) indicates clear diffraction spots arising from the fine particles of Si phase that precipitated within the α-Al phases. A scanningTEM (STEM) image and the corresponding EDS elemental mapping images are presented in Fig. 8. The present EDS chemical analysis revealed the columnar α-Al phase surrounded by fine Si particles (Fig. 8(a–c)). Minute point analyses revealed that a certain amount of Si

at 530 °C, relatively coarse particles of Si phase with a mean size above 2 μm were often observed in the α-Al matrix (Fig. 4(e)), which corresponds well to the formation of the indistinct melt-pool boundaries observed in the optical micrographs (Fig. 2(e)). Significantly coarsened particles of Si phase and a rod-shaped Fe-rich intermetallic phase (βAlFeSi) were observed in the α-Al matrix after long-term annealing (Fig. 4(f)). In order to quantify the changes in the area fractions of the particles of Si phase (fSi) dispersed in the α-Al matrix during annealing at different temperatures, area fractions for all of the observed samples were measured by carrying out image analysis on several SEM micrographs. Note that extremely fine particles with a size of smaller than 10 nm (below the resolution level of SEM) were excluded from the obtained results. The quantified values of fSi are summarized in Fig. 6, together with the calculated fSi values in equilibrium measured from thermodynamic calculations [23]. The present calculations for the Al–Si–Mg–Fe quaternary system were performed using a thermodynamic database for an Al-based multi-component system (PanAluminum) [29] for the measured alloy composition of Al–10.8Si–0.2Mg–0.4Fe (wt%), as listed in Table 1. The average fSi value experimentally measured in the as-built sample was 10.3 %. During annealing at 300 °C, the fSi value slightly increased up until 2 h and then reached 12.3 % after 6 h. The fSi value appeared to be saturated after annealing for 6 h, which corresponds well to the calculated fSi value at equilibrium at 300 °C. The fSi value continuously increased upon an increase in the annealing time at 530 °C. The fSi increased to approximately 11.5 % (which is almost same as the calculated value in 5

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Fig. 6. Change in the area fraction of the Si particles, fSi (measured from SEM images) for different annealing times at different temperatures of 300 and 530 °C.

α-Al grains in the high-magnification TEM image (Fig. 9(c)). Note that the TEM image was captured after electron beam irradiation for a few minutes, which is indicative of a slight effect of electron beam damage on the observed microstructure. These puncta contribute to the inhomogeneous contrast variation observed in the microstructure of the as-fabricated sample (Fig. 7(a)). These puncta have been previously observed in the α-Al phase in the SLM-fabricated samples of different Al–Si-based alloys [12,24,26,30]. A HRTEM image showing a region containing a nanoscale contrast is presented in Fig. 9(d). Fast Fourier transform (FFT) treatments of the HRTEM image were performed, as shown in Fig. 9(e). The FFT image represents the incident beam along the [0 1 1] direction of the α-Al (fcc) matrix. It is noteworthy that no clear reflections were detected, even in the selected area containing a nano-sized contrast, indicating that the observed numerous puncta might not be derived from any of the precipitates (e.g., the Si or Mg2Si phases) within the α-Al phase. By selecting the pair of reflection vectors of g1-11, g11-1, and g200, an inverse FFT was performed to show any dislocations associated with the observed dot contrast, the results of which are shown in Fig. 9(f–h). IFFT analyses revealed that a few dislocations were detected inside/around the dot in the α-Al (fcc) matrix, which is indicative of a few lattice defects associated with the observed numerous puncta. These results suggest that the observed contrast variations might be responsible for the formation of solution clusters (likely with fcc structure) or their related phases in the α-Al matrix, which is in good agreement with the 3D atom probe analysis results for the α-Al phase in the SLM-built AlSi10Mg alloy [30]. Note that these puncta were rarely observed in the annealed sample (Fig. 7(c)), which corresponds well to the previous observations [32].

Fig. 5. High-magnification SEM images showing fine Si particles in (a) as-fabricated AlSi10Mg alloy sample and (b, c) subsequently annealed samples: (b) 300 °C for 60 s, (c) 300 °C for 2 h.

was detected inside the columnar α-Al phase (Fig. 8(c)), whereas the measured Si concentration were scattered from 0.6 to 2.2 at%, indicating the difficulty in precise quantification of solute Si content in the α-Al phase using the present STEM-EDS analyses. The high Si content detected is consistent with the previous results of various SLMbuilt Al–Si-based alloys [12,24,26,30]. Note that Fe was found to be segregated in the α-Al/Si eutectic microstructure (Fig. 8(d)), whereas detectable enriched Mg was not found (Fig. 8(e)). In order to identify the inhomogeneous contrast variations observed within the α-Al phase in the as-fabricated sample, HRTEM observations were carried out and the results can be seen in Fig. 9. A TEM-bright field image and the corresponding SAED pattern obtained from the location inside the columnar α-Al grain (Fig. 9(a, b)) indicate the incident beam parallel to the [0 1 1] zone axis of the α-Al (fcc) phase. The SAED pattern (Fig. 9(b)) shows that no extra reflections were derived from the Si (diamond structure) or Mg2Si [31] phases, whereas numerous nanoscale puncta (with dark contrast) are apparent within the

3.2. Thermal conductivity Fig. 10 shows the thermal conductivities of the as-built sample and subsequently annealed samples measured at ambient temperature. The thermal conductivities perpendicular and parallel to the building direction (Z direction) were designated as λX/Y and λZ, respectively. The measured λX/Y and λZ values of the as-fabricated sample are 97 and 108 W·m−1 K−1, respectively, which correspond well to the reported material data sheet of the SLM-built AlSi10Mg alloy [21]. The difference between λX/Y and λZ is due to the morphologies of the elongated αAl grains surrounded by the Si particles, as presented in Fig. 4(a). It is noteworthy that the thermal conductivity increases upon annealing at elevated temperatures. The measured thermal conductivities (λX/Y and λZ) of the 300 °C/60 s annealed sample are approximately 155 W·m−1 K−1. The measured λX/Y and λZ values slightly increased to approximately 162 W·m−1 K−1 after annealing for 2 h. The thermal 6

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Fig. 7. (a, c) TEM images and (b, d) selected area electron diffraction patterns for (a, b) as-fabricated AlSi10Mg alloy sample and (c, d) subsequently annealed sample (300 °C for 2 h).

Fig. 8. (a) STEM images of the as-built AlSi10Mg alloy sample and the (d-e) corresponding EDS elemental mapping images of (b) Al, (c) Fe, (d) Si, and (e) Mg.

In order to investigate the changes in the strength of the as-built specimen, tensile tests were carried out on the as-built and annealed samples. The results are summarized in Fig. 12. Flow curves (Fig. 12(a)) were obtained from the tensile-deformation perpendicular to the building direction (along the X/Y direction). The true stress (σ)–true strain (ε) curve of the as-built specimen indicated a yield strength of approximately 250 MPa. The flow stress increased at a high strain hardening rate (dσ /dε) and then reached above 500 MPa. The 300 °C annealed specimen exhibited a lower yield strength. The following strain hardening rate was obviously lower than that of the as-fabricated sample, resulting in a lower tensile strength. This trend was also observed for the 530 °C annealed specimen. The changes in the strain hardening rate (dσ /dε) of the tested specimens were plotted as a function of the true strain (ε), together with their true stress–strain (σ–ε) curves according to the Considère criterion for plastic instability

conductivities of the 530 °C/6 h annealed sample are approximately 153 W·m−1 K−1, independent of the measurement direction (λX/Y and λZ). The isotropic thermal conductivity is due to the presence of the homogenously coarsened Si particles in the α-Al matrix (Fig. 4(f)). 3.3. Hardness and strength Fig. 11 shows the changes in the hardness of the as-built sample according to the annealing time at different temperatures. The as-built sample exhibits a high hardness of 131 HV. After annealing at 300 °C, the hardness significantly decreased to 105 HV, even after 60 s. The hardness continuously decreased to 92 HV after 2 h and then reached 83 HV after 6 h. At a higher temperature of 530 °C, the hardness dropped to approximately 85 HV after 60 s and then became almost saturated after further annealing. 7

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Fig. 9. (a) TEM-bright field image (BFI) showing the microstructure of the as-built AlSi10Mg alloy sample, (b) SAED pattern of the columnar α-Al grain, (c) highmagnification BFI, (d) HRTEM image of the nanoscale particles, (e) the corresponding fast Fourier-transformed (FFT) image and (f-h) inverse FFT (IFFT) images using different reciprocal lattice vectors (b).

Fig. 10. Thermal conductivities of the as-built sample and subsequently annealed samples measured using a laser flash method at room temperature.

Fig. 11. Change in the hardness of the SLM-built AlSi10Mg alloy sample with annealing time at different temperatures of 300 and 530 °C.

8

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precipitate within the α-Al matrix. These results indicate that the fine Si precipitate within the α-Al phase played no role in the strengthening of the AlSi10Mg alloy. The systematic analyses for the area fraction of the Si particles using the observed SEM images (Fig. 6) demonstrated that a certain amount of solute Si in the α-Al matrix was present in the as-built sample, indicating that a supersaturated solid solution of the α-Al phase formed due to the rapid solidification (at an extremely high cooling rate of about 105 K·s−1 [35]) during the SLM process. The concentrated solute atoms in the α-Al matrix were supported by the low thermal conductivity of the as-built sample comparing to those of the annealed samples (Fig. 10). At elevated temperatures, the Si phase (with a diamond structure) precipitated consuming the solute Si in the α-Al matrix (Fig. 7(c)). These experimental results show that the reduced hardness (strength) upon annealing is due to a loss of solute Si in the α-Al matrix. The present results suggest that the concentrated solute Si in the αAl matrix would be a dominant contributing factor in strengthening of the SLM-fabricated AlSi10Mg alloy. One of the most important findings for understanding the strength is the observation of nano-sized particles in the α-Al matrix of the as-built sample (Fig. 9). This observation is consistent with a previous atom-probe tomography results that indicated the presence of atom-concentrated clusters in the SLM-built AlSi10Mg alloy [30]. Furthermore, the puncta contrast derived from the nano-sized particles has been observed in various Al–Si-based alloys fabricated by SLM [12,24,26,30]. These nanoscale particles cannot be detected at the SEM resolution level (Fig. 5(a)) and appear as a concentrated Si solute in the α-Al phase when evaluated by image analyses using the SEM images (Fig. 6). It is generally known that in conventional Al alloys, nanoscale clusters significantly contribute toward the strengthening of heat-treated alloys [36], whereby it can be assumed that the observed nano-sized particles play the role of acting as obstacles for the dislocation motion in the α-Al matrix, resulting in the strengthening of the SLM-built AlSi10Mg alloy. This assumption is in reasonable agreement with the high strain hardening rate observed in the as-built specimen during the tensile deformation (Fig. 12(b)). The nano-sized particles may be able to limit the dislocation movement, resulting in enhanced strain hardening. In order to confirm the proposed mechanism, the dislocation substructures in the tensile deformed samples were observed by TEM. The resulting bright-field images showing the microstructures of the as-built and 300 °C annealed specimens tensile deformed at 3% strain are displayed in Fig. 13. In the asbuilt tensile deformed sample, an extremely high density of dislocations were observed within the columnar α-Al phase in the cellular eutectic microstructure (Fig. 13(a)). The accumulated dislocations were observed around the cell boundaries surrounding the α-Al phase as well. These dislocation lines appear to significantly interact with each other (Fig. 13(c)) and their high density is in good agreement with the high internal stress enhanced by strain hardening. This result corresponds well to the previous result of a high density of dislocations inside the columnar α-Al phase introduced by TEM during in situ loading [20]. A number of dislocations were also observed in the 300 °C annealed specimen tensile-deformed (Fig. 13(b)). The dislocations often appear to interact with the fine Si precipitate in the α-Al matrix (Fig. 13(d)), whereas the morphologies imply only a slight interaction among the dislocations. It is obvious that their density was much lower than that observed in the as-built sample (Fig. 13(a)), indicating that the fine Si precipitate dispersed in the α-Al phase (detectable at the SEM resolution level) slightly contributes toward the limiting of the dislocation motion. These observations are consistent with the difference in the strain hardening rate (Fig. 13(b)), which supports our proposed mechanism for strengthening by the supersaturated solid solution of the αAl phase (including nano-sized particles) in the SLM-built AlSi10Mg alloy. Based on the aforementioned discussion, the present results could provide new insights in understanding the anomalous strengthening of Al alloys produced by SLM, as presented in Fig. 1. The as-fabricated AlSi10Mg alloy exhibits columnar α-Al phase (a primary solidified

Fig. 12. (a) True stress-strain curves and (b) strain hardening rates of the asbuilt AlSi10Mg alloy and subsequently annealed alloy specimens measured by tensile tests at room temperature.

in rate-insensitive materials (σ > dσ/dε) [33], the summarized results of which are shown in Fig. 11(b). The as-built specimen exhibits a high strain hardening rate (dσ/dε) at an early stage of tensile deformation (below 0.07 in strain). It is noteworthy that the dσ/dε values of the asbuilt specimen are more than twice those of the annealed specimens. The reduction in the dσ/dε does not coincide with the flow stress (σ) in the as-built specimen, which corresponds well to the slight macroscopic necking of the tensile-fractured specimen of the as-built AlSi10Mg alloy [34]. The dσ/dε values of the 300 °C annealed specimen were found to be almost the same as those of the 530 °C annealed specimen at a lower strain of below 0.05. This result obviously indicates that the fine Si phase precipitated within the α-Al phase (Figs. 5(c) and 7(c)) slightly contributes toward the strain hardening and its associated tensile strength. 4. Discussion The present study set out to examine the changes in the strength (hardness) and the associated microstructural parameters of an SLMbuilt AlSi10Mg alloy upon annealing at different temperatures. A fine particle Si phase was precipitated within the columnar α-Al phase surrounded by a eutectic Si phase at 300 °C (Figs. 5(b, c), 7(c)), whereas the microstructure of the α-Al matrix slightly changed upon annealing (Fig. 3). It is noteworthy that the hardness substantially decreased upon annealing at 300 °C (Fig. 11), regardless of the presence of the fine Si 9

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Fig. 13. TEM bright-field images showing dislocation substructures inside the columnar α-Al phase in the cellar microstructures of (a,c) as-fabricated specimen and (b,d) 300 °C/ 2 h annealed specimen tensile-deformed to 3% strain at room temperature.

it can be concluded that the supersaturated solid solution of the α-Al phase (including numerous nano-sized particles) is a dominant contributor to the anomalous strengthening of Al–Si-based alloys produced by SLM. However, there is still an issue regarding the identification of the detailed structure and constitute elements of the nano-sized particles observed in the SLM-fabricated AlSi10Mg alloy. The formation of the supersaturated solid solution of the α-Al phase could be formed during rapid solidification, whereas the solidification microstructure might be heated by additional laser irradiation on the upper powder layer in the SLM process [1,2]. The effect of the additional heating process on the formation of nano-sized particles and its association with the strength of the SLM-fabricated Al alloys remains unclear. Furthermore, heating base plates (approximately 200 °C) are often used to reduce the thermal gradient and stress of the fabricated samples. The heating plate may reduce the cooling rate of the solidified samples, resulting in the fine Si precipitates (dissolution of nano-sized particles) [38–40]. Actually, the AlSi10Mg alloy sample fabricated by SLM using the heating base plate exhibited a relatively low strength (approximately 350∼400 MPa) [18,19]. Thus, further microstructural characterization of the Al alloy samples prepared under various laser conditions (laser powder or scanning speed) and atmosphere temperatures is required to allow the controlling of the strength level of the SLM-fabricated Al–Si-based alloys.

phase) surrounded by eutectic Si phases, which developed as a result of rapid solidification at an extremely high cooling rate in the SLM process. The rapidly solidified α-Al phase contains a concentrated Si in solution and these solute atoms may form nano-sized particles that disperse within the α-Al phase. These particles could limit the dislocation motion, resulting in enhanced strain hardening. The observed nano-sized particles (Fig. 9) appear somewhat larger than nanoscale clusters generally found in age-hardened Al alloys [36], whereby these particles might play a similar role as precipitation hardening rather than strengthening by atomic clusters. The proposed mechanism suggests a slight strengthening of Al alloys containing low amounts of alloy elements by SLM. This suggestion is in good agreement with the slight difference in strength observed between SLM-fabricated and conventionally cast pure Al (Fig. 1(a)). It has been reported that fine Si phase precipitates in the series of Al–Si-based alloys aged at relatively low temperatures [12,24,26], whereby the strengthening by SLM would be more significant for Al–Si-based alloys with a high Si content (Fig. 1(a)). The summarized data imply that the solubility limit of Si might determine the strength of the SLM-fabricated Al–Si-based alloys. It is intriguing that T6 heat treatments (solution treatment and subsequent artificial aging) significantly reduce the strengthening effect by the SLM process (Fig. 1(b)). During high temperature (500∼540 °C) solution treatments, coarsening of the Si phase and the formation of a stable intermetallic phase occur (Fig. 4(e,f)), resulting in a reduction in the solute Si in the α-Al matrix (and nano-sized particles may dissolve into the α-Al matrix). In fact, nano-sized particles and their associated varied contrasts were not observed in the annealed sample (Fig. 7(c)). Thus, the age hardening of the SLM-fabricated solution-treated Al alloys could be similar to that of the conventionally cast alloys. The aforementioned mechanism is consistent with the recent study [37] on strengthening by nano-precipitates of the SLM-fabricated Al alloy treated by direct aging (without the solution treatment). Consequently,

5. Summary The present study was performed to identify the dominant microstructural features responsible for the high strength of an AlSi10Mg alloy built using SLM. The results of the investigation regarding the supersaturated solid solutions of the α-Al phase were used to understand the changes in the microstructure at elevated temperatures, as 10

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well as their associated mechanical and thermal properties. The following conclusions were shown from the results and discussion.

[10]

(1) Microstructure consisting of a number of columnar α-Al phase with concentrated Si in solution was observed in the as-built AlSi10Mg alloy. In the α-Al matrix, numerous nano-sized particles were observed. At elevated temperatures, the Si phase finely precipitated within the columnar α-Al grains, but the α-Al matrix microstructure slightly changed. The consumption of the solute Si in the α-Al matrix by the precipitated Si phase was supported by the much lower thermal conductivity of the as-fabricated AlSi10Mg alloy than those of the annealed samples. (2) The hardness of as-fabricated AlSi10Mg alloy substantially decreased upon annealing at elevated temperatures. The reduced hardness (strength) corresponds well to the reduced strain hardening rate after annealing. (3) The supersaturated solid solution of the α-Al phase (including nanosized particles) formed due to rapid solidification at an extremely high cooling rate is a dominant contributor to the anomalous strengthening of AlSi10Mg alloys produced by SLM. The present results provide important insights into controlling the strength of Al–Si-based alloys fabricated by SLM process.

[11]

[12]

[13]

[14]

[15] [16]

[17] [18]

[19]

[20]

Authorship statements

[21]

N. Takata: Conception and design by study; acquisition of data; analysis and/or interpretation of data; Drafting the manuscript; revising the manuscript for important intellectual content. A. Suzuki: Conception and design by study; revising the manuscript for important intellectual content. M. Kobashi: Conception and design by study; revising the manuscript for important intellectual content. M. Liu: Acquistion of data; analysis and/or interpretation of data. H. kodaira: Acquisition of data.

[22] [23]

[24]

[25]

[26]

Declaration of Competing Interest None.

[27]

Acknowledgments

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The supports of the “Knowledge Hub Aichi”, a Priority Research Project of Aichi Prefectural Government, Japan and the JSPS KAKENHI (grant number 17H03411) are gratefully acknowledged. We are grateful for on sample preparation provided by Mr. H. Tauchi of WhiteImapct Co., Ltd.

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[30] [31] [32]

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