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Effect of bimodal microstructure on the tensile properties of selective laser melt Al-Mg-Sc-Zr alloy
Journal of Alloys and Compounds 815 (2020) 152422
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of bimodal microstructure on the tensile properties of selective laser melt Al-Mg-Sc-Zr alloy Rulong Ma a, Chaoqun Peng a, Zhiyong Cai a, b, *, Richu Wang a, b, Zhaohui Zhou c, Xiaogeng Li c, Xuanyang Cao c a
School of Materials Science and Engineering, Central South University, Changsha, 410083, China National Key Laboratory of Science and Technology for National Defence on High-strength Structural Materials, Central South University, Changsha, 410083, China c Changsha Advanced Materials Industrial Research Institute, Changsha, 410083, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 June 2019 Received in revised form 22 September 2019 Accepted 23 September 2019 Available online 24 September 2019
Selective laser melting (SLM) has become a promising manufacturing route for the aluminum alloy models and parts in automotive and aircraft industries. In this work, an Al-4.0Mg-0.7Sc-0.4Zr-0.5Mn alloy was fabricated by SLM and near-full densification (relative density 99.7%) was obtained. A fine equiaxed grain regions and columnar grain regions alternately distributed characteristic is observed in the as-fabricated sample. The results of microstructural analyses show that the rapid solidification process and the precipitation of Al3Mg2, Al6(Fe, Mn), and Al3(Sc, Zr) particles lead to the fine grain structure, which provides a combination of high yield strength and excellent ductility. The mechanical properties of the as-fabricated sample were analyzed based on the typical microstructural feature of a combination of coarse columnar grain regions and fine grain regions. The strengthening effect is mainly attributed to the grain boundaries and the effect of grain boundaries on yield stress is according with the classical Hall-Petch relationship. © 2019 Elsevier B.V. All rights reserved.
Keywords: Selective laser melting Al-Mg alloy Bimodal microstructure Tensile properties Strengthening mechanism
1. Introduction Aluminum alloys are increasingly employed in automotive and aircraft applications because of their low density, high strength to weight ratio, low cost, and excellent formability. Aluminum alloy parts have been prepared by conventional manufacturing methods such as casting, forging, extrusion, and powder metallurgy. However, these conventional manufacturing methods have disadvantages of time-consuming and expensive processing, especially for parts with complex geometry. In recent years, great interest has been devoted to additive manufacturing (AM) due to its enormous advantages, such as short design and production period, high material utilization ratio, time saving, and the ability to prepare complex parts. Selective laser melting (SLM), one type of AM, has become a promising manufacturing route for metallic models and parts. SLM creates parts by layer-by-layer fabrication through the application of laser energy to powder bed based on the three-
* Corresponding author. School of Materials Science and Engineering, Central South University, Changsha, 410083, China. E-mail address: [email protected] (Z. Cai). https://doi.org/10.1016/j.jallcom.2019.152422 0925-8388/© 2019 Elsevier B.V. All rights reserved.
dimensional (3D) CAD data. Fine grain structures with excellent mechanical performance can be obtained due to rapid heating and high solidification rates (up to 105e106 K/s) in the SLM process. At present, SLM has been applied extensively in aluminum alloys, but this application is mainly limited in Al-Si alloys such as Al-10Si-Mg [1,2], Al-12Si [3e8], and Al-7Si-Mg [9]. However, there has been a few works on Al-Mg alloys by SLM. These alloys have extensive prospects as structural materials in the above sectors due to their superior strength, high ductility, and low anisotropic mechanical behavior in the as-fabricated condition. Spierings et al. [10] reported that an SLM scandium-modified aluminum alloy exhibits alternately fine-grained microstructure and coarser grains. Mechanical properties of this alloy overcome the values reported for Al-Si alloys, with tensile strength Rm value exceeding 500 MPa and show a low anisotropy with regard to the building orientation [11]. These properties have been attributed to the ultra-fine grain structure, along with the excellent hardenability of the alloy. For SLM Al-6.2Mg-0.36Sc-0.09Zr, the maximal compression strength (390.25 MPa) and electrochemical corrosion property are higher than those of the cast one [12]. The effects of post-processing heat treatment on the misconstruction and
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2. Experimental procedures
Table 1 Chemical composition of the Al-Mg-Sc-Zr alloy (wt.%). Element
Mg
Mn
Sc
Zr
Fe
Si
Al
Content
3.72
0.57
0.74
0.42
0.15
0.059
Bal.
mechanical properties of the Al-Mg-Sc-Zr alloy were further studied [11]. Bimodal grain size distribution microstructure was observed in the SLM processed aluminum alloys. This is a typical feature that is different from other traditional preparation methods. An understanding of how the coarse columnar grain (CG) regions and fine grain (FG) regions affect the physical and mechanical properties is essential for the design of the SLM aluminum alloys. However, the microstructure difference and effect on the tensile properties between the CG and the FG regions have not been carefully studied. In this work, we investigated the difference of the FG and CG regions and its influence on the mechanical properties of the AlMg-Sc-Zr alloy in the as-built condition. The strengthening mechanism of the SLM alloy was also discussed.
2.1. Raw materials The Al-Mg-Sc-Zr powder was produced by inert gas atomization with a nominal composition of Al-4.0Mg-0.7Sc-0.4Zr-0.5Mn. The raw materials were made of high-purity aluminum ingot (>99.9 wt %), pure magnesium (>99.9 wt%), Al-10Mn (>99.9 wt%), Al-10Zr (>99.9 wt%), and Al-2Sc master alloys. Atomization was carried out at 1.5e5 MPa and 780e820 C with an Ar gas atomizer. The entire atomization process was carried out under Ar gas protection. The chemical composition of the alloy powder was determined by inductively coupled plasma optical emission spectroscopy (ICPOES), and the result is listed in Table 1. The aluminum and magnesium contents of the powder almost correspond to the target composition. The scanning electron microscopy (SEM) images of the morphology, microstructure, and size distribution of the gasatomized alloy powder are shown in Fig. 1. The powder particles are almost spherical in shape and contain a large number of satellites. The microstructure of the atomized powder is dendritic as shown in the upper right corner of Fig. 1(a). The mechanically sieved powder particle size ranges from 7 to 68 mm, with a mean particle diameter (d50) of 24.3 mm.
Fig. 1. SEM images of (a) morphology and (b) size distribution of the gas-atomized Al-Mg-Sc-Zr alloy powder.
Fig. 2. (a) Schematic illustration of the SLM process and (b) the shape and size of as-SLM samples.
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Fig. 3. The specimen design scheme for tensile tests.
2.2. SLM process The experiments were carried out on a selective laser melting instrument (AM400, Renishaw) equipped with an Ytterbium fiber laser of 400 W class (laser focus diameter: approximately 70 mm). The samples were fabricated using a laser power of 325 W, a scan speed of 1200 mm/s, and a hatch spacing of 80 mm under an Ar atmosphere with less than 0.1% oxygen content. The slice thickness was kept to 30 mm. The volumetric energy density (E) is given as
E¼
P v,t,d
(1)
where P is the laser power (W), v is the scan speed (mm/s), t is the slice thickness (mm), and d is the hatch spacing (mm). According to the formula, the applied E is 113 J/mm3. Finally, 10 10 10 mm3 cube samples were prepared for the analysis of density, microhardness, and microstructure, whereas 10 10 100 mm3 samples were prepared for static tensile test, as shown in Fig. 2. Additionally, the cast alloy containing the same composition was also prepared for comparison.
2.3. Property measurements and characterization Sectional planes parallel to the building direction of the SLM samples were mirror-finished by cutting, grinding, polishing, and etched using Keller’s reagent. The microstructures were then observed using optical microscopy (OM, MA100L, Nikon), field emission scanning electron microscopy (SEM, SIRION200, FEI) equipped with an energy-dispersive X-ray spectroscopy (EDX) setup and double beam electron microscope instrument (SEM, Helios Nanolab G3 UC, FEI) equipped with electron backscatter diffraction (EBSD). Phase analysis was carried out by X-ray diffraction (XRD) using a Rigaku D/max instrument with Cu Ka radiation (0.154054 nm) at a scan step of 8 ( )/min. Element
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distribution was measured using a JXA-8230 electron probe microanalyzer (EPMA). The TEM samples preparation consisted of cutting slices 0.8 mm in thickness by electric spark cutting, reducing their thickness with SiC sand paper and 0.5 mm diamond suspensions down to about 80 mm, punching out of them 3 mm discs and electropolishing the discs in 30 vol% Nitric Acid solution in Methanol operated at 17 V. TEM analysis was performed in a Tecnai G20ST microscope operated at 200 kV. Both bright field (BF) and high angle annular dark field (HAADF) images were captured in scanning transmission electron microscopy (STEM) mode. Image J software was adopted to calculate the content and average size of the fine and coarse grains based on the OM and TEM images, respectively. The size of the grains and precipitates was measured using at least ten randomly selected HAADF-STEM micrographs. The grain size applied is the average of the values estimated by TEM and EBSD. The density of the fabricated samples was measured by the Archimedes method. The microhardness was measured with a Shimadzu HMV-2T microhardness meter. The tensile samples were designed according to ASTM E8M, the cross-section being 3 mm thick, 6 mm wide in the gauge section, and 100 mm in length, as shown in Fig. 3. At least three tensile samples were tested at room temperature using a MTS810 testing machine under quasistatic loading at a tensile rate of 2 mm/min and the strain was measured directly on the specimens using a Fiedler laser-extensometer under each condition. The fracture surfaces of the specimens were characterized by SEM using a SIRION200 microscope. 3. Results and discussion 3.1. Microstructure and precipitates The microstructure of the SLM sample produced with the building direction alternating from layer to layer is shown in Fig. 4(a). Only a few homogenous and small micropores are found on the surface in Fig. 4(a). There are no noticeable large pores. Generally, the SLM sample is basically dense, and the relative density of the SLM sample reaches 99.7%. The porosity can be improved by increasing the sphericity of the material powder and adjusting the processing parameters. The alternate-layered microstructure of the Al-Mg-Sc-Zr alloy along the building direction is shown in the upper right corner of Fig. 4(a). Overlapped melt pools can be clearly observed. From the weld-lines being prepared perpendicular to the building direction, the typical thickness of each layer is approximately 30 mm. Fig. 4(b) shows the microstructure of the cast sample. As can be seen, the microstructure is uniform and the grain size is in the range of 50e90 mm.
Fig. 4. Microstructures of (a) the SLM sample and (b) the cast sample.
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Fig. 5. Backscattered SEM images of the SLM Al-Mg-Sc-Zr alloy with (a) low magnification and (b) high magnification.
Fig. 5 shows the backscattered SEM images of the SLM Al-MgSc-Zr alloy. Two different regions can be distinguished clearly: the FG regions with a circular equiaxed grain structure and the CG regions which exhibit elongated and enlarged columnar grains. The two regions are alternately distributed. This phenomenon illustrates that the solidification process is dominated by the thermal gradient. Compared with the traditional casting method, the SLM samples have different structural and mechanical properties due to the use of a laser as a heating source. During the SLM process, the cooling rate near the surface of the melt pool area reached
1 103e1.44 106 K/s [8,13,14]. Therefore, fine grains can be obtained by rapid solidification. The overlap areas between two adjacent melt pool tracks will be melted for two or more times, depending on the scan strategy and processing parameters. During solidification, the grains tend to grow along the direction of the positive temperature gradient from the remelted zone to the top of the melt pool. Therefore, the columnar grains form and grow toward the center of the melt pool. Fig. 6(a) shows the image quality (IQ) mapping result and Fig. 6(b) presents the inverse pole figure (IPF) mapping. As can be
Fig. 6. (a) IQ mapping and (b) IPF mapping of the sample, showing coarse and fine grained regions. (001) pole figures for (c) FG and (d) CG regions.
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Fig. 7. Grain size distribution in (a) FG and (b) CG regions and grain orientation angle distribution in (c) FG and (d) CG regions.
Fig. 8. TEM bright-field images of (a) coarse grains and (b) fine grains of the SLM Al-Mg-Sc-Zr alloy.
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R. Ma et al. / Journal of Alloys and Compounds 815 (2020) 152422 Table 2 Microstructure characteristics of the SLM Al-Mg-Sc-Zr alloy. Parameter
Method
Fine grain
Coarse grain
Grain diameter (mm)
EBSD TEM EBSD EBSD TEM EBSD TEM OM Microhardness tester
0.60 ± 0.26 0.50 ± 0.20 19.7 e e e e 35 101 ± 3
3.5 ± 1.8 e 32.0 6.0 ± 4.0 5.5 ± 3.3 2.0 ± 0.92 2.4 ± 0.80 65 95 ± 3
Average misorientation angle ( ) Grain length (mm) Grain width (mm) Content (%) Microhardness (Hv)
Fig. 9. XRD analysis of the SLM and the cast Al-Mg-Sc-Zr alloys.
seen, the microstructure consists of fine and coarse grains regions, which is consistent with the metallographic structure. In the FG regions, the color difference of adjacent grains is larger, which indicates different orientations. However, the orientation of adjacent grains in the CG regions is not much different. As shown in Fig. 6(c) and (d), the (001) pole figures for the FG and CG regions also show random grain orientation in the fine grain regions, and a preferential grain orientation in the coarse grain regions. A <100> texture is found in the CG regions with a maximum texture intensity of around 7.2, which is typical for SLM-processed materials. The grain growth of CG regions is perpendicular to the melt pool boundary due to the dominance of the temperature gradient, indicating that different solidification behaviors for FG and CG regions. The grain size and grain orientation angle distribution in the FG and CG regions are shown in Fig. 7. As can be seen, the grain size is less than 2.5 mm and randomly oriented in the FG regions. However, the grain size in the CG regions is up to 6 mm, which is still smaller than that of the cast alloy. It can be seen from Fig. 7(c) and (d) that there is a large amount of low angle grain boundaries in the FG regions. Nevertheless, there are almost no low angle grain boundaries in the CG regions. Fig. 8 shows the microstructure observed by TEM. The microstructure characteristics of the SLM Al-Mg-Sc-Zr alloy are listed in Table 2. Obvious grain boundaries and second phases can be seen in Fig. 8. Bimodal grain size distribution is observed in the SLM sample, including coarse columnar grains (with an average width of 2.2 ± 0.92 mm and length of 5.5 ± 4.0 mm) and fine equiaxed grains (with a size of 550 ± 0.20 nm). The microhardness values of the FG
and CG regions are 102 Hv and 95 Hv, respectively. XRD investigations were carried out to identify the phases present in the samples, and the results are illustrated in Fig. 9. It can be seen that the diffraction peaks of Al, Al3Mg2, and Al6Mn phases are observed in the cast and SLM Al-Mg-Sc-Zr alloys. Combined with SEM þ EDS and XRD results, it can be concluded that the AlMg-Sc-Zr alloy contains Al, Al3Mg2, and Al6(Fe, Mn) phases. Compared with the as-cast sample, the diffraction peaks of the SLM sample are broader, indicating considerable grain refinement and super-saturation caused by laser rapid melting/solidification. This result is consistent with the reports on Al-Mg [15,16] and Al-Si alloys [6,17]. EPMA was used to analyze the distribution of the main alloy elements, and the results are illustrated in Fig. 10. It can be observed that Al, Mn, Sc, and Zr elements are basically homogeneously distributed in the alloy. The main reason for this phenomenon is that Mn, Sc and Zr are mainly dissolved in the aluminum matrix due to the high cooling rate, and only a little residual is mainly present in the Al-containing compounds which are formed at the beginning of solidification and cannot be detected in low magnifications. It is obviously demonstrated that there is a certain degree of composition segregation of Mg, as illustrated in Fig. 10(c). Fig. 11 shows the HAADF-TEM images of the SLM Al-Mg-Sc-Zr alloy. Relatively large bright particles in Fig. 11 are enriched in high atomic number elements (i.e. Mg, Mn and Fe) and are mainly distributed along the grain boundaries. This is consistent with the results of other traditional Al-Mg alloys. Mg and Mn are more likely to precipitate at grain boundaries. Small spherical phases are uniformly distributed in the grain interiors and their average size is 20e40 nm. The small phases at a low content are mainly primary phases, and their main role is to refine the grains during the solidification process, but the dispersion strengthening effect is weak. With the aim of gaining insights into the distribution of the principal alloying elements, TEM-EDX analysis was conducted. The bright field image (Fig. 11 frame selection area) and its corresponding marked region EDX maps of Mg, Mn, Sc, Zr and Fe are shown in Fig. 12. The large particles (point 1) contain Mg as the major element, indicating the Al3Mg2 precipitates, which mainly nucleate along the grain boundaries. Chemical analysis of the relatively large precipitates (point 2) with a rod-like shape reveals a high amount of Mn and Fe, and these particles are identified as Al6(Fe, Mn) precipitates. Nanosized phases within the grains consist of Sc or Zr, indicating the Al3(Sc, Zr) precipitates. Mg and Mn composition segregations in the Al3(Sc, Zr) precipitates are found in Fig. 12. When Mg is added to the Al alloys, it dissolves into the a-Al phase resulting in solid solution strengthening. However, the Al3Mg2 phase is mainly precipitated along the grain boundaries during the cooling process, so its aging strengthening effect does not have much practical value. Al3Mg2 has no momentous strengthening effect and reduces the corrosion response [18]. Mn
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Fig. 10. EPMA element maps of the SLM Al-Mg-Sc-Zr alloy with the observation plane parallel to the building direction.
Fig. 11. HAADF-TEM images of the SLM Al-Mg-Sc-Zr alloy: (a) CG regions and (b) FG regions.
element mainly acts as additional strengthening additive. During the solidification process, the Al6Mn phase can also dissolve a part of impurity element Fe to form an Al6(Fe, Mn) phase, which reduces the harmful effects of Fe. Similar to the Al3Mg2 phase, the Al6Mn
phase precipitates from the supersaturated solid solution is also mainly distributed on the grain boundaries. Sc is commonly used as a grain-refining element in conventional cast Al alloys. Zr is used to replace part of Sc to improve the thermal stability of the precipitate
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Fig. 12. TEMeEDX analysis of the SLM Al-Mg-Sc-Zr alloy. (a)Bright field image and corresponding marked region EDX maps of (b) Mg, (c) Mn, (d) Sc, (e) Zr and (f) Fe.
Zr) precipitate strengthening. Sc and Zr are used as grain refining agents by forming primary Al3Sc or Al3(Sc, Zr) phases in the melt. These precipitates can act as the nucleation centers of Al grains and produce significant grain refinement. Additionally, this mechanism can explain partially the cause of forming the bimodal grain size distribution. Slow solidification front velocities allow for more primary Al3Sc precipitation at the edge of the melt pool, which leads to fine grains with random orientation. The grain growth is inhibited by the Zener pinning due to the Al3Sc precipitation in the FG materials [21e23] and the Mg segregation at the grain boundaries (GBs). Because of the high solidification-front velocity at the top of the melt pool, Sc and Zr elements are more effectively trapped in the Al matrix, which leads to little primary Al3Sc precipitation. Fig. 11(a) shows that the coarse-grained regions have fewer Al3Sc particles than the fine-grained regions. The coarse grains are nucleated by the underlying fine grains and solidify directionally along the temperature gradient. 3.2. Mechanical characterization Fig. 13. Tensile curves of the SLM and cast Al-Mg-Sc-Zr alloys at room temperature.
Table 3 Room temperature tensile properties of the as-SLM and as-cast Al-Mg-Sc-Zr alloys. Sample
su (MPa)
sy (MPa)
ε (%)
E (GPa)
as-SLM as-cast
362 ± 7 264 ± 5
332 ± 10 198 ± 8
11.2 ± 2 25.0 ± 2
69.1 ± 1.4 19.1 ± 1.6
phase and reduce the cost. Maximum solid solubility of Sc and Zr in solid Al alloys are 0.35 wt% at 665 C and 0.083 wt% at 660 C [19,20]. The strengthening mechanisms of the Al-Mg-Sc-Zr alloy include solid solution strengthening, grain refine strengthening, and Al3(Sc,
Fig. 13 shows the measured stress-strain curves of the SLM and cast Al-Mg-Sc-Zr alloys. The tensile properties are summarized in Table 3. Compared with the cast alloy, the SLM alloy has relatively higher tensile strength and lower elongation, which are 362 ± 7 MPa and 11.2 ± 2%, respectively. Additionally, a gentle strain hardening stage is observed in the SLM alloy, indicating a weak strain hardening behavior. The work hardening effect is weak on account of the accumulation of limited dislocations and storage within the grains, which is consistent with the results of fine grained Al-Mg alloys [24]. The as-SLM alloy has a high elongation compared to un-bimodal fine-grained materials due to the bimodal grain size distribution. On account of lack of work hardening capability, ultrafine-grained metal materials generally reveal limited uniform elongation, usually less than 5% in tension. According to the model put forward by
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Fig. 14. SEM images showing the characteristic fracture surfaces (a), (b), (c) and cross-sectional images near the fracture (c) CG þ FG regions, (d) FG regions, (e) CG regions of the SLM Al-Mg-Sc-Zr sample.
Raeisinia et al. [25], the bimodal grain size distribution can improve the tensile elongation and provide better overall properties compared with the un-bimodal counterparts. The coarse grains can accommodate more dislocations and elongation and result in a high tensile ductility. The superior strength is mainly derived from the fine grains which are not easily deformed and can withstand more stress than the coarse grains. It has been verified that ultrafinegrained and fine-grained alloys with a bimodal microstructure exhibit a combination of high strength and good ductility [26e29]. The morphologies of the tensile fracture surfaces are shown in Fig. 14. As can be seen in Fig. 14(a), most of the sample fracture surfaces are at an angle of 45 to the direction of stretching, that is to say, the sample fractured in the direction of maximum shear stresses. For the as-SLM alloy, small dimples are found on the fracture surfaces indicating a certain degree of ductility. Brittle characteristics appear on the fracture surfaces perpendicular to the applied load (parallel to the building direction) at the center, which can be verified by the lamellar and transgranular fractures in Fig. 14(c). This phenomenon indicates a weak bonding force between the coarse grains. Elongated grains near the fracture surface can be observed in Fig. 14(e) and (f). The coarse grains themselves are involved in the plastic deformation, thereby delaying the crack initiation and propagation. Except for the dispersion strengthening effect provided by the Al3Sc precipitates and the solid solution strengthening effect provided by the Mg atoms in the lattice, the grain boundaries can
effectively hinder the dislocation motion. The influence of grain size on yield strength can be predicted by the Hall-Petch relationship,
sy ¼ s0 þ kd1=2
(2)
where sy is the yield stress, s0 is the friction stress, k is an experimental constant, and d is the average grain diameter. The k value ranges from 0.15 to 0.20 MPa m1/2 for Al-(3e6)%Mg alloy [30e33]. The Hall-Petch relationship has been confirmed for a lot of materials. We wonder the applicability of the relationship for the Al-Mg alloy processed by SLM and the results have been verified by experiments. For grain size with d ¼ 3.5 mm and d ¼ 550 nm, Eq. (2) predicts an increase in yield strength of DsGB ¼ 90 MPa and DsGB ¼ 230 MPa, respectively, when k is 0.17. The average strengthening provided by the bimodal grain size is expected to be DsGB ¼ 230 0.35 þ 90 0.65 ¼ 139 MPa. This simple estimation ignores the connectivity and interactions of the coarse and fine grains regions. Additionally, it assumes a straightforward linear superposition to estimate the grain boundary strengthening by the HallePetch mechanism of the alloy. The yield strength is 332 MPa for the as-SLM Al-Mg-Sc-Zr alloy, whereas it is 198 MPa for the ascast Al-Mg-Sc-Zr alloy. The strengthening is Ds ¼ 332e198 ¼ 134 MPa. This is in good agreement with the estimated value by the Hall-Petch relationship. This result indicates that the strengthening effect of SLM alloy is mainly attributed to the grain boundaries.
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4. Conclusion In this work, an Al-Mg-Sc-Zr alloy was successfully manufactured by SLM. The microstructure and mechanical properties were investigated, and their relationship was discussed. The following conclusions can be drawn: (1) Dense samples with a 99.7% relative density, 362 ± 7 MPa tensile strength, and 11.2% elongation are obtained in the asfabricated alloys at a laser power density of 113 J/mm3. (2) The microstructure of the SLM alloy is composed of coarse columnar grains (with an average width of 2.2 ± 0.90 mm and length of 6.0 ± 4.0 mm) separated by fine equiaxed grains (with size of 0.55 ± 0.26 mm). The strengthening effect is mainly attributed to the grain boundaries and the effect of grain boundaries on yield stress accords with the Hall-Petch relationship. (3) The Al3Mg2 and Al6Mn precipitates are mainly distributed along the grain boundaries, while primary Al3(Sc, Zr) particles act as a grain refining agent and are mainly distributed within the grains. Acknowledgements The authors are grateful for the financial support from the National Natural Science Foundation of China (51804349), the China Postdoctoral Science Foundation (2018M632986), the Natural Science Foundation of Hunan Province (2019JJ50766), the Postdoctoral Science Foundation of Central South University, and the Science and Technology Plan Projects of Hunan Province (2017GK2261). References [1] M. Tang, Inclusions, Porosity, and Fatigue of AlSi10Mg Parts Produced by Selective Laser Melting, Carnegie Mellon University, Pittsburgh, 2017. [2] U. Tradowsky, J. White, R.M. Ward, N. Read, W. Reimers, M.M. Attallah, Selective laser melting of AlSi10Mg: influence of post-processing on the microstructural and tensile properties development, Mater. Des. 105 (2016) 212e222. [3] K.G. Prashanth, J. Eckert, Formation of metastable cellular microstructures in selective laser melted alloys, J. Alloy. Comp. 707 (2017) 27e34. [4] T. Kimura, T. Nakamoto, M. Mizuno, H. Araki, Effect of silicon content on densification, mechanical and thermal properties of Al-xSi binary alloys fabricated using selective laser melting, Mater. Sci. Eng., A 682 (2017) 593e602. [5] X.P. Li, K.M. O’Donnell, T.B. Sercombe, Selective laser melting of Al-12Si alloy: enhanced densification via powder drying, Addit. Manuf. 10 (2016) 10e14. [6] K.G. Prashanth, S. Scudino, J. Eckert, Defining the tensile properties of Al-12Si parts produced by selective laser melting, Acta Mater. 126 (2017) 25e35. [7] N. Kang, P. Coddet, L. Dembinski, H.L. Liao, C. Coddet, Microstructure and strength analysis of eutectic Al-Si alloy in-situ manufactured using selective laser melting from elemental powder mixture, J. Alloy. Comp. 691 (2017) 316e322. [8] J. Suryawanshi, K.G. Prashanth, S. Scudino, J. Eckert, O. Prakash, U. Ramamurty, Simultaneous enhancements of strength and toughness in an Al-12Si alloy synthesized using selective laser melting, Acta Mater. 115 (2016) 285e294. [9] K.V. Yang, P. Rometsch, C.H.J. Davies, A.J. Huang, X.H. Wu, Effect of heat treatment on the microstructure and anisotropy in mechanical properties of A357 alloy produced by selective laser melting, Mater. Des. 154 (2018)
275e290. [10] A.B. Spierings, K. Dawson, M. Voegtlin, F. Palm, P.J. Uggowitzer, Microstructure and mechanical properties of as-processed scandium-modified aluminium using selective laser melting, CIRP Ann 65 (2016) 213e216. [11] A.B. Spierings, K. Dawson, K. Kern, F. Palm, K. Wegener, SLM-processed Scand Zr- modified Al-Mg alloy: mechanical properties and microstructural effects of heat treatment, Mater. Sci. Eng., A 701 (2017) 264e273. [12] R.D. Li, M.B. Wang, T.C. Yuan, B. Song, C. Chen, K.C. Zhou, P. Cao, Selective laser melting of a novel Sc and Zr modified Al-6.2 Mg alloy: processing, microstructure, and properties, Powder Technol. 319 (2017) 117e128. [13] P. Wang, H.C. Li, K.G. Prashanth, J. Eckert, S. Scudino, Selective laser melting of Al-Zn-Mg-Cu: heat treatment, microstructure and mechanical properties, J. Alloy. Comp. 707 (2017) 287e290. [14] E.O. Olakanmi, R.F. Cochrane, K.W. Dalgarno, A review on selective laser sintering/melting (SLS/SLM) of aluminium alloy powders: processing, microstructure, and properties, Prog. Mater. Sci. 74 (2015) 401e477. [15] A.B. Spierings, K. Dawson, P. Dumitraschkewitz, S. Pogatscher, K. Wegener, Microstructure characterization of SLM-processed Al-Mg-Sc-Zr alloy in the heat treated and HIPed condition, Addit. Manuf. 20 (2018) 173e181. [16] J.R. Croteau, S. Griffiths, M.D. Rossell, C. Leinenbach, C. Kenel, V. Jansen, D.N. Seidman, D.C. Dunand, N.Q. Vo, Microstructure and mechanical properties of Al-Mg-Zr alloys processed by selective laser melting, Acta Mater. 153 (2018). [17] Y.J. Liu, Z. Liu, Y. Jiang, G.W. Wang, Y. Yang, L.C. Zhang, Gradient in microstructure and mechanical property of selective laser melted AlSi10Mg, J. Alloy. Comp. 735 (2018) 1414e1421. [18] G.E. Totten, D.S. Mackenzie, Handbook of Aluminum, Crc Press, 2003. [19] J. Murray, A. Peruzzi, J.P. Abriata, The Al-Zr (aluminum-zirconium) system, J. Phase Equilibria 13 (1992) 277e291. [20] J.L. Murray, The Al-Sc (aluminum-scandium) system, J. Phase Equilibria 19 (1998) 380e384. [21] K. Ma, E.J. Lavernia, J.M. Schoenung, Absorption of nitrogen at Al/Al2O3 interfaces in Al nanocomposites: a computational analysis, Adv. Eng. Mater. 14 (2012) 77e84. [22] M.A. Kaka, M. Julie Schoenung, Influence of cryomilling on the microstructural features in HVOF-sprayed NiCrAIY bond coats for thermal barrier coatings: creation of a homogeneous distribution of nanoscale dispersoids, Philos. Mag. Lett. 90 (2010) 739e751. [23] V.L. Tellkamp, E.J. Lavernia, A. Melmed, Mechanical behavior and microstructure of a thermally stable bulk nanostructured Al alloy, Metall. Mater. Trans. A 32 (2001) 2335e2343. [24] T.J. Harrell, T.D. Topping, H. Wen, T. Hu, J.M. Schoenung, E.J. Lavernia, Microstructure and strengthening mechanisms in an ultrafine grained Al-MgSc alloy produced by powder metallurgy, Metall. Mater. Trans. A 45 (2014) 6329e6343. , On the impact of grain size [25] B. Raeisinia, C.W. Sinclair, W.J. Poole, C.N. Tome distribution on the plastic behaviour of polycrystalline metals, Model. Simul. Mater. Sci. Eng. 16 (2008), 025001. [26] D. Witkin, Z. Lee, R. Rodriguez, S. Nutt, E. Lavernia, AleMg alloy engineered with bimodal grain size for high strength and increased ductility, Scr. Mater. 49 (2003) 297e302. [27] G.J. Fan, H. Choo, P.K. Liaw, E.J. Lavernia, Plastic deformation and fracture of ultrafine-grained AleMg alloys with a bimodal grain size distribution, Acta Mater. 54 (2006) 1759e1766. [28] M.S. Oskooie, H. Asgharzadeh, H.S. Kim, Microstructure, plastic deformation and strengthening mechanisms of an Al-Mg-Si alloy with a bimodal grain structure, J. Alloy. Comp. 632 (2015) 540e548. [29] Z. Lee, V. Radmilovic, B. Ahn, E.J. Lavernia, S.R. Nutt, Tensile deformation and fracture mechanism of bulk bimodal ultrafine-grained Al-Mg alloy, Metall. Mater. Trans. A 41 (2010) 795e801. [30] H. Hasegawa, S. Komura, A. Utsunomiya, Z. Horita, M. Furukawa, M. Nemoto, T.G. Langdon, Thermal stability of ultrafine-grained aluminum in the presence of Mg and Zr additions, Mater. Sci. Eng., A 265 (1999) 188e196. [31] E.L. Huskins, B. Cao, K.T. Ramesh, Strengthening mechanisms in an AleMg alloy, Mater. Sci. Eng., A 527 (2010) 1292e1298. [32] A. Dubyna, A. Mogucheva, R. Kaibyshev, Hall-petch relationship in an Al-MgSc alloy subjected to ECAP, Adv. Mater. Res. 922 (2014) 120e125. [33] D.J. Lloyd, S.A. Court, Influence of grain size on tensile properties of Al-Mg alloys, Met. Sci. J. 19 (2014) 1349e1354.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of bimodal microstructure on the tensile properties of selective laser melt Al-Mg-Sc-Zr alloy Rulong Ma a, Chaoqun Peng a, Zhiyong Cai a, b, *, Richu Wang a, b, Zhaohui Zhou c, Xiaogeng Li c, Xuanyang Cao c a
School of Materials Science and Engineering, Central South University, Changsha, 410083, China National Key Laboratory of Science and Technology for National Defence on High-strength Structural Materials, Central South University, Changsha, 410083, China c Changsha Advanced Materials Industrial Research Institute, Changsha, 410083, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 June 2019 Received in revised form 22 September 2019 Accepted 23 September 2019 Available online 24 September 2019
Selective laser melting (SLM) has become a promising manufacturing route for the aluminum alloy models and parts in automotive and aircraft industries. In this work, an Al-4.0Mg-0.7Sc-0.4Zr-0.5Mn alloy was fabricated by SLM and near-full densification (relative density 99.7%) was obtained. A fine equiaxed grain regions and columnar grain regions alternately distributed characteristic is observed in the as-fabricated sample. The results of microstructural analyses show that the rapid solidification process and the precipitation of Al3Mg2, Al6(Fe, Mn), and Al3(Sc, Zr) particles lead to the fine grain structure, which provides a combination of high yield strength and excellent ductility. The mechanical properties of the as-fabricated sample were analyzed based on the typical microstructural feature of a combination of coarse columnar grain regions and fine grain regions. The strengthening effect is mainly attributed to the grain boundaries and the effect of grain boundaries on yield stress is according with the classical Hall-Petch relationship. © 2019 Elsevier B.V. All rights reserved.
Keywords: Selective laser melting Al-Mg alloy Bimodal microstructure Tensile properties Strengthening mechanism
1. Introduction Aluminum alloys are increasingly employed in automotive and aircraft applications because of their low density, high strength to weight ratio, low cost, and excellent formability. Aluminum alloy parts have been prepared by conventional manufacturing methods such as casting, forging, extrusion, and powder metallurgy. However, these conventional manufacturing methods have disadvantages of time-consuming and expensive processing, especially for parts with complex geometry. In recent years, great interest has been devoted to additive manufacturing (AM) due to its enormous advantages, such as short design and production period, high material utilization ratio, time saving, and the ability to prepare complex parts. Selective laser melting (SLM), one type of AM, has become a promising manufacturing route for metallic models and parts. SLM creates parts by layer-by-layer fabrication through the application of laser energy to powder bed based on the three-
* Corresponding author. School of Materials Science and Engineering, Central South University, Changsha, 410083, China. E-mail address: [email protected] (Z. Cai). https://doi.org/10.1016/j.jallcom.2019.152422 0925-8388/© 2019 Elsevier B.V. All rights reserved.
dimensional (3D) CAD data. Fine grain structures with excellent mechanical performance can be obtained due to rapid heating and high solidification rates (up to 105e106 K/s) in the SLM process. At present, SLM has been applied extensively in aluminum alloys, but this application is mainly limited in Al-Si alloys such as Al-10Si-Mg [1,2], Al-12Si [3e8], and Al-7Si-Mg [9]. However, there has been a few works on Al-Mg alloys by SLM. These alloys have extensive prospects as structural materials in the above sectors due to their superior strength, high ductility, and low anisotropic mechanical behavior in the as-fabricated condition. Spierings et al. [10] reported that an SLM scandium-modified aluminum alloy exhibits alternately fine-grained microstructure and coarser grains. Mechanical properties of this alloy overcome the values reported for Al-Si alloys, with tensile strength Rm value exceeding 500 MPa and show a low anisotropy with regard to the building orientation [11]. These properties have been attributed to the ultra-fine grain structure, along with the excellent hardenability of the alloy. For SLM Al-6.2Mg-0.36Sc-0.09Zr, the maximal compression strength (390.25 MPa) and electrochemical corrosion property are higher than those of the cast one [12]. The effects of post-processing heat treatment on the misconstruction and
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2. Experimental procedures
Table 1 Chemical composition of the Al-Mg-Sc-Zr alloy (wt.%). Element
Mg
Mn
Sc
Zr
Fe
Si
Al
Content
3.72
0.57
0.74
0.42
0.15
0.059
Bal.
mechanical properties of the Al-Mg-Sc-Zr alloy were further studied [11]. Bimodal grain size distribution microstructure was observed in the SLM processed aluminum alloys. This is a typical feature that is different from other traditional preparation methods. An understanding of how the coarse columnar grain (CG) regions and fine grain (FG) regions affect the physical and mechanical properties is essential for the design of the SLM aluminum alloys. However, the microstructure difference and effect on the tensile properties between the CG and the FG regions have not been carefully studied. In this work, we investigated the difference of the FG and CG regions and its influence on the mechanical properties of the AlMg-Sc-Zr alloy in the as-built condition. The strengthening mechanism of the SLM alloy was also discussed.
2.1. Raw materials The Al-Mg-Sc-Zr powder was produced by inert gas atomization with a nominal composition of Al-4.0Mg-0.7Sc-0.4Zr-0.5Mn. The raw materials were made of high-purity aluminum ingot (>99.9 wt %), pure magnesium (>99.9 wt%), Al-10Mn (>99.9 wt%), Al-10Zr (>99.9 wt%), and Al-2Sc master alloys. Atomization was carried out at 1.5e5 MPa and 780e820 C with an Ar gas atomizer. The entire atomization process was carried out under Ar gas protection. The chemical composition of the alloy powder was determined by inductively coupled plasma optical emission spectroscopy (ICPOES), and the result is listed in Table 1. The aluminum and magnesium contents of the powder almost correspond to the target composition. The scanning electron microscopy (SEM) images of the morphology, microstructure, and size distribution of the gasatomized alloy powder are shown in Fig. 1. The powder particles are almost spherical in shape and contain a large number of satellites. The microstructure of the atomized powder is dendritic as shown in the upper right corner of Fig. 1(a). The mechanically sieved powder particle size ranges from 7 to 68 mm, with a mean particle diameter (d50) of 24.3 mm.
Fig. 1. SEM images of (a) morphology and (b) size distribution of the gas-atomized Al-Mg-Sc-Zr alloy powder.
Fig. 2. (a) Schematic illustration of the SLM process and (b) the shape and size of as-SLM samples.
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Fig. 3. The specimen design scheme for tensile tests.
2.2. SLM process The experiments were carried out on a selective laser melting instrument (AM400, Renishaw) equipped with an Ytterbium fiber laser of 400 W class (laser focus diameter: approximately 70 mm). The samples were fabricated using a laser power of 325 W, a scan speed of 1200 mm/s, and a hatch spacing of 80 mm under an Ar atmosphere with less than 0.1% oxygen content. The slice thickness was kept to 30 mm. The volumetric energy density (E) is given as
E¼
P v,t,d
(1)
where P is the laser power (W), v is the scan speed (mm/s), t is the slice thickness (mm), and d is the hatch spacing (mm). According to the formula, the applied E is 113 J/mm3. Finally, 10 10 10 mm3 cube samples were prepared for the analysis of density, microhardness, and microstructure, whereas 10 10 100 mm3 samples were prepared for static tensile test, as shown in Fig. 2. Additionally, the cast alloy containing the same composition was also prepared for comparison.
2.3. Property measurements and characterization Sectional planes parallel to the building direction of the SLM samples were mirror-finished by cutting, grinding, polishing, and etched using Keller’s reagent. The microstructures were then observed using optical microscopy (OM, MA100L, Nikon), field emission scanning electron microscopy (SEM, SIRION200, FEI) equipped with an energy-dispersive X-ray spectroscopy (EDX) setup and double beam electron microscope instrument (SEM, Helios Nanolab G3 UC, FEI) equipped with electron backscatter diffraction (EBSD). Phase analysis was carried out by X-ray diffraction (XRD) using a Rigaku D/max instrument with Cu Ka radiation (0.154054 nm) at a scan step of 8 ( )/min. Element
3
distribution was measured using a JXA-8230 electron probe microanalyzer (EPMA). The TEM samples preparation consisted of cutting slices 0.8 mm in thickness by electric spark cutting, reducing their thickness with SiC sand paper and 0.5 mm diamond suspensions down to about 80 mm, punching out of them 3 mm discs and electropolishing the discs in 30 vol% Nitric Acid solution in Methanol operated at 17 V. TEM analysis was performed in a Tecnai G20ST microscope operated at 200 kV. Both bright field (BF) and high angle annular dark field (HAADF) images were captured in scanning transmission electron microscopy (STEM) mode. Image J software was adopted to calculate the content and average size of the fine and coarse grains based on the OM and TEM images, respectively. The size of the grains and precipitates was measured using at least ten randomly selected HAADF-STEM micrographs. The grain size applied is the average of the values estimated by TEM and EBSD. The density of the fabricated samples was measured by the Archimedes method. The microhardness was measured with a Shimadzu HMV-2T microhardness meter. The tensile samples were designed according to ASTM E8M, the cross-section being 3 mm thick, 6 mm wide in the gauge section, and 100 mm in length, as shown in Fig. 3. At least three tensile samples were tested at room temperature using a MTS810 testing machine under quasistatic loading at a tensile rate of 2 mm/min and the strain was measured directly on the specimens using a Fiedler laser-extensometer under each condition. The fracture surfaces of the specimens were characterized by SEM using a SIRION200 microscope. 3. Results and discussion 3.1. Microstructure and precipitates The microstructure of the SLM sample produced with the building direction alternating from layer to layer is shown in Fig. 4(a). Only a few homogenous and small micropores are found on the surface in Fig. 4(a). There are no noticeable large pores. Generally, the SLM sample is basically dense, and the relative density of the SLM sample reaches 99.7%. The porosity can be improved by increasing the sphericity of the material powder and adjusting the processing parameters. The alternate-layered microstructure of the Al-Mg-Sc-Zr alloy along the building direction is shown in the upper right corner of Fig. 4(a). Overlapped melt pools can be clearly observed. From the weld-lines being prepared perpendicular to the building direction, the typical thickness of each layer is approximately 30 mm. Fig. 4(b) shows the microstructure of the cast sample. As can be seen, the microstructure is uniform and the grain size is in the range of 50e90 mm.
Fig. 4. Microstructures of (a) the SLM sample and (b) the cast sample.
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Fig. 5. Backscattered SEM images of the SLM Al-Mg-Sc-Zr alloy with (a) low magnification and (b) high magnification.
Fig. 5 shows the backscattered SEM images of the SLM Al-MgSc-Zr alloy. Two different regions can be distinguished clearly: the FG regions with a circular equiaxed grain structure and the CG regions which exhibit elongated and enlarged columnar grains. The two regions are alternately distributed. This phenomenon illustrates that the solidification process is dominated by the thermal gradient. Compared with the traditional casting method, the SLM samples have different structural and mechanical properties due to the use of a laser as a heating source. During the SLM process, the cooling rate near the surface of the melt pool area reached
1 103e1.44 106 K/s [8,13,14]. Therefore, fine grains can be obtained by rapid solidification. The overlap areas between two adjacent melt pool tracks will be melted for two or more times, depending on the scan strategy and processing parameters. During solidification, the grains tend to grow along the direction of the positive temperature gradient from the remelted zone to the top of the melt pool. Therefore, the columnar grains form and grow toward the center of the melt pool. Fig. 6(a) shows the image quality (IQ) mapping result and Fig. 6(b) presents the inverse pole figure (IPF) mapping. As can be
Fig. 6. (a) IQ mapping and (b) IPF mapping of the sample, showing coarse and fine grained regions. (001) pole figures for (c) FG and (d) CG regions.
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Fig. 7. Grain size distribution in (a) FG and (b) CG regions and grain orientation angle distribution in (c) FG and (d) CG regions.
Fig. 8. TEM bright-field images of (a) coarse grains and (b) fine grains of the SLM Al-Mg-Sc-Zr alloy.
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R. Ma et al. / Journal of Alloys and Compounds 815 (2020) 152422 Table 2 Microstructure characteristics of the SLM Al-Mg-Sc-Zr alloy. Parameter
Method
Fine grain
Coarse grain
Grain diameter (mm)
EBSD TEM EBSD EBSD TEM EBSD TEM OM Microhardness tester
0.60 ± 0.26 0.50 ± 0.20 19.7 e e e e 35 101 ± 3
3.5 ± 1.8 e 32.0 6.0 ± 4.0 5.5 ± 3.3 2.0 ± 0.92 2.4 ± 0.80 65 95 ± 3
Average misorientation angle ( ) Grain length (mm) Grain width (mm) Content (%) Microhardness (Hv)
Fig. 9. XRD analysis of the SLM and the cast Al-Mg-Sc-Zr alloys.
seen, the microstructure consists of fine and coarse grains regions, which is consistent with the metallographic structure. In the FG regions, the color difference of adjacent grains is larger, which indicates different orientations. However, the orientation of adjacent grains in the CG regions is not much different. As shown in Fig. 6(c) and (d), the (001) pole figures for the FG and CG regions also show random grain orientation in the fine grain regions, and a preferential grain orientation in the coarse grain regions. A <100> texture is found in the CG regions with a maximum texture intensity of around 7.2, which is typical for SLM-processed materials. The grain growth of CG regions is perpendicular to the melt pool boundary due to the dominance of the temperature gradient, indicating that different solidification behaviors for FG and CG regions. The grain size and grain orientation angle distribution in the FG and CG regions are shown in Fig. 7. As can be seen, the grain size is less than 2.5 mm and randomly oriented in the FG regions. However, the grain size in the CG regions is up to 6 mm, which is still smaller than that of the cast alloy. It can be seen from Fig. 7(c) and (d) that there is a large amount of low angle grain boundaries in the FG regions. Nevertheless, there are almost no low angle grain boundaries in the CG regions. Fig. 8 shows the microstructure observed by TEM. The microstructure characteristics of the SLM Al-Mg-Sc-Zr alloy are listed in Table 2. Obvious grain boundaries and second phases can be seen in Fig. 8. Bimodal grain size distribution is observed in the SLM sample, including coarse columnar grains (with an average width of 2.2 ± 0.92 mm and length of 5.5 ± 4.0 mm) and fine equiaxed grains (with a size of 550 ± 0.20 nm). The microhardness values of the FG
and CG regions are 102 Hv and 95 Hv, respectively. XRD investigations were carried out to identify the phases present in the samples, and the results are illustrated in Fig. 9. It can be seen that the diffraction peaks of Al, Al3Mg2, and Al6Mn phases are observed in the cast and SLM Al-Mg-Sc-Zr alloys. Combined with SEM þ EDS and XRD results, it can be concluded that the AlMg-Sc-Zr alloy contains Al, Al3Mg2, and Al6(Fe, Mn) phases. Compared with the as-cast sample, the diffraction peaks of the SLM sample are broader, indicating considerable grain refinement and super-saturation caused by laser rapid melting/solidification. This result is consistent with the reports on Al-Mg [15,16] and Al-Si alloys [6,17]. EPMA was used to analyze the distribution of the main alloy elements, and the results are illustrated in Fig. 10. It can be observed that Al, Mn, Sc, and Zr elements are basically homogeneously distributed in the alloy. The main reason for this phenomenon is that Mn, Sc and Zr are mainly dissolved in the aluminum matrix due to the high cooling rate, and only a little residual is mainly present in the Al-containing compounds which are formed at the beginning of solidification and cannot be detected in low magnifications. It is obviously demonstrated that there is a certain degree of composition segregation of Mg, as illustrated in Fig. 10(c). Fig. 11 shows the HAADF-TEM images of the SLM Al-Mg-Sc-Zr alloy. Relatively large bright particles in Fig. 11 are enriched in high atomic number elements (i.e. Mg, Mn and Fe) and are mainly distributed along the grain boundaries. This is consistent with the results of other traditional Al-Mg alloys. Mg and Mn are more likely to precipitate at grain boundaries. Small spherical phases are uniformly distributed in the grain interiors and their average size is 20e40 nm. The small phases at a low content are mainly primary phases, and their main role is to refine the grains during the solidification process, but the dispersion strengthening effect is weak. With the aim of gaining insights into the distribution of the principal alloying elements, TEM-EDX analysis was conducted. The bright field image (Fig. 11 frame selection area) and its corresponding marked region EDX maps of Mg, Mn, Sc, Zr and Fe are shown in Fig. 12. The large particles (point 1) contain Mg as the major element, indicating the Al3Mg2 precipitates, which mainly nucleate along the grain boundaries. Chemical analysis of the relatively large precipitates (point 2) with a rod-like shape reveals a high amount of Mn and Fe, and these particles are identified as Al6(Fe, Mn) precipitates. Nanosized phases within the grains consist of Sc or Zr, indicating the Al3(Sc, Zr) precipitates. Mg and Mn composition segregations in the Al3(Sc, Zr) precipitates are found in Fig. 12. When Mg is added to the Al alloys, it dissolves into the a-Al phase resulting in solid solution strengthening. However, the Al3Mg2 phase is mainly precipitated along the grain boundaries during the cooling process, so its aging strengthening effect does not have much practical value. Al3Mg2 has no momentous strengthening effect and reduces the corrosion response [18]. Mn
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Fig. 10. EPMA element maps of the SLM Al-Mg-Sc-Zr alloy with the observation plane parallel to the building direction.
Fig. 11. HAADF-TEM images of the SLM Al-Mg-Sc-Zr alloy: (a) CG regions and (b) FG regions.
element mainly acts as additional strengthening additive. During the solidification process, the Al6Mn phase can also dissolve a part of impurity element Fe to form an Al6(Fe, Mn) phase, which reduces the harmful effects of Fe. Similar to the Al3Mg2 phase, the Al6Mn
phase precipitates from the supersaturated solid solution is also mainly distributed on the grain boundaries. Sc is commonly used as a grain-refining element in conventional cast Al alloys. Zr is used to replace part of Sc to improve the thermal stability of the precipitate
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Fig. 12. TEMeEDX analysis of the SLM Al-Mg-Sc-Zr alloy. (a)Bright field image and corresponding marked region EDX maps of (b) Mg, (c) Mn, (d) Sc, (e) Zr and (f) Fe.
Zr) precipitate strengthening. Sc and Zr are used as grain refining agents by forming primary Al3Sc or Al3(Sc, Zr) phases in the melt. These precipitates can act as the nucleation centers of Al grains and produce significant grain refinement. Additionally, this mechanism can explain partially the cause of forming the bimodal grain size distribution. Slow solidification front velocities allow for more primary Al3Sc precipitation at the edge of the melt pool, which leads to fine grains with random orientation. The grain growth is inhibited by the Zener pinning due to the Al3Sc precipitation in the FG materials [21e23] and the Mg segregation at the grain boundaries (GBs). Because of the high solidification-front velocity at the top of the melt pool, Sc and Zr elements are more effectively trapped in the Al matrix, which leads to little primary Al3Sc precipitation. Fig. 11(a) shows that the coarse-grained regions have fewer Al3Sc particles than the fine-grained regions. The coarse grains are nucleated by the underlying fine grains and solidify directionally along the temperature gradient. 3.2. Mechanical characterization Fig. 13. Tensile curves of the SLM and cast Al-Mg-Sc-Zr alloys at room temperature.
Table 3 Room temperature tensile properties of the as-SLM and as-cast Al-Mg-Sc-Zr alloys. Sample
su (MPa)
sy (MPa)
ε (%)
E (GPa)
as-SLM as-cast
362 ± 7 264 ± 5
332 ± 10 198 ± 8
11.2 ± 2 25.0 ± 2
69.1 ± 1.4 19.1 ± 1.6
phase and reduce the cost. Maximum solid solubility of Sc and Zr in solid Al alloys are 0.35 wt% at 665 C and 0.083 wt% at 660 C [19,20]. The strengthening mechanisms of the Al-Mg-Sc-Zr alloy include solid solution strengthening, grain refine strengthening, and Al3(Sc,
Fig. 13 shows the measured stress-strain curves of the SLM and cast Al-Mg-Sc-Zr alloys. The tensile properties are summarized in Table 3. Compared with the cast alloy, the SLM alloy has relatively higher tensile strength and lower elongation, which are 362 ± 7 MPa and 11.2 ± 2%, respectively. Additionally, a gentle strain hardening stage is observed in the SLM alloy, indicating a weak strain hardening behavior. The work hardening effect is weak on account of the accumulation of limited dislocations and storage within the grains, which is consistent with the results of fine grained Al-Mg alloys [24]. The as-SLM alloy has a high elongation compared to un-bimodal fine-grained materials due to the bimodal grain size distribution. On account of lack of work hardening capability, ultrafine-grained metal materials generally reveal limited uniform elongation, usually less than 5% in tension. According to the model put forward by
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Fig. 14. SEM images showing the characteristic fracture surfaces (a), (b), (c) and cross-sectional images near the fracture (c) CG þ FG regions, (d) FG regions, (e) CG regions of the SLM Al-Mg-Sc-Zr sample.
Raeisinia et al. [25], the bimodal grain size distribution can improve the tensile elongation and provide better overall properties compared with the un-bimodal counterparts. The coarse grains can accommodate more dislocations and elongation and result in a high tensile ductility. The superior strength is mainly derived from the fine grains which are not easily deformed and can withstand more stress than the coarse grains. It has been verified that ultrafinegrained and fine-grained alloys with a bimodal microstructure exhibit a combination of high strength and good ductility [26e29]. The morphologies of the tensile fracture surfaces are shown in Fig. 14. As can be seen in Fig. 14(a), most of the sample fracture surfaces are at an angle of 45 to the direction of stretching, that is to say, the sample fractured in the direction of maximum shear stresses. For the as-SLM alloy, small dimples are found on the fracture surfaces indicating a certain degree of ductility. Brittle characteristics appear on the fracture surfaces perpendicular to the applied load (parallel to the building direction) at the center, which can be verified by the lamellar and transgranular fractures in Fig. 14(c). This phenomenon indicates a weak bonding force between the coarse grains. Elongated grains near the fracture surface can be observed in Fig. 14(e) and (f). The coarse grains themselves are involved in the plastic deformation, thereby delaying the crack initiation and propagation. Except for the dispersion strengthening effect provided by the Al3Sc precipitates and the solid solution strengthening effect provided by the Mg atoms in the lattice, the grain boundaries can
effectively hinder the dislocation motion. The influence of grain size on yield strength can be predicted by the Hall-Petch relationship,
sy ¼ s0 þ kd1=2
(2)
where sy is the yield stress, s0 is the friction stress, k is an experimental constant, and d is the average grain diameter. The k value ranges from 0.15 to 0.20 MPa m1/2 for Al-(3e6)%Mg alloy [30e33]. The Hall-Petch relationship has been confirmed for a lot of materials. We wonder the applicability of the relationship for the Al-Mg alloy processed by SLM and the results have been verified by experiments. For grain size with d ¼ 3.5 mm and d ¼ 550 nm, Eq. (2) predicts an increase in yield strength of DsGB ¼ 90 MPa and DsGB ¼ 230 MPa, respectively, when k is 0.17. The average strengthening provided by the bimodal grain size is expected to be DsGB ¼ 230 0.35 þ 90 0.65 ¼ 139 MPa. This simple estimation ignores the connectivity and interactions of the coarse and fine grains regions. Additionally, it assumes a straightforward linear superposition to estimate the grain boundary strengthening by the HallePetch mechanism of the alloy. The yield strength is 332 MPa for the as-SLM Al-Mg-Sc-Zr alloy, whereas it is 198 MPa for the ascast Al-Mg-Sc-Zr alloy. The strengthening is Ds ¼ 332e198 ¼ 134 MPa. This is in good agreement with the estimated value by the Hall-Petch relationship. This result indicates that the strengthening effect of SLM alloy is mainly attributed to the grain boundaries.
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4. Conclusion In this work, an Al-Mg-Sc-Zr alloy was successfully manufactured by SLM. The microstructure and mechanical properties were investigated, and their relationship was discussed. The following conclusions can be drawn: (1) Dense samples with a 99.7% relative density, 362 ± 7 MPa tensile strength, and 11.2% elongation are obtained in the asfabricated alloys at a laser power density of 113 J/mm3. (2) The microstructure of the SLM alloy is composed of coarse columnar grains (with an average width of 2.2 ± 0.90 mm and length of 6.0 ± 4.0 mm) separated by fine equiaxed grains (with size of 0.55 ± 0.26 mm). The strengthening effect is mainly attributed to the grain boundaries and the effect of grain boundaries on yield stress accords with the Hall-Petch relationship. (3) The Al3Mg2 and Al6Mn precipitates are mainly distributed along the grain boundaries, while primary Al3(Sc, Zr) particles act as a grain refining agent and are mainly distributed within the grains. Acknowledgements The authors are grateful for the financial support from the National Natural Science Foundation of China (51804349), the China Postdoctoral Science Foundation (2018M632986), the Natural Science Foundation of Hunan Province (2019JJ50766), the Postdoctoral Science Foundation of Central South University, and the Science and Technology Plan Projects of Hunan Province (2017GK2261). References [1] M. Tang, Inclusions, Porosity, and Fatigue of AlSi10Mg Parts Produced by Selective Laser Melting, Carnegie Mellon University, Pittsburgh, 2017. [2] U. Tradowsky, J. White, R.M. Ward, N. Read, W. Reimers, M.M. Attallah, Selective laser melting of AlSi10Mg: influence of post-processing on the microstructural and tensile properties development, Mater. Des. 105 (2016) 212e222. [3] K.G. Prashanth, J. Eckert, Formation of metastable cellular microstructures in selective laser melted alloys, J. Alloy. Comp. 707 (2017) 27e34. [4] T. Kimura, T. Nakamoto, M. Mizuno, H. Araki, Effect of silicon content on densification, mechanical and thermal properties of Al-xSi binary alloys fabricated using selective laser melting, Mater. Sci. Eng., A 682 (2017) 593e602. [5] X.P. Li, K.M. O’Donnell, T.B. Sercombe, Selective laser melting of Al-12Si alloy: enhanced densification via powder drying, Addit. Manuf. 10 (2016) 10e14. [6] K.G. Prashanth, S. Scudino, J. Eckert, Defining the tensile properties of Al-12Si parts produced by selective laser melting, Acta Mater. 126 (2017) 25e35. [7] N. Kang, P. Coddet, L. Dembinski, H.L. Liao, C. Coddet, Microstructure and strength analysis of eutectic Al-Si alloy in-situ manufactured using selective laser melting from elemental powder mixture, J. Alloy. Comp. 691 (2017) 316e322. [8] J. Suryawanshi, K.G. Prashanth, S. Scudino, J. Eckert, O. Prakash, U. Ramamurty, Simultaneous enhancements of strength and toughness in an Al-12Si alloy synthesized using selective laser melting, Acta Mater. 115 (2016) 285e294. [9] K.V. Yang, P. Rometsch, C.H.J. Davies, A.J. Huang, X.H. Wu, Effect of heat treatment on the microstructure and anisotropy in mechanical properties of A357 alloy produced by selective laser melting, Mater. Des. 154 (2018)
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