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Microstructure and mechanical properties of a novel Sc and Zr modified 7075 aluminum alloy prepared by selective laser melting
Materials Science & Engineering A 768 (2019) 138478
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Microstructure and mechanical properties of a novel Sc and Zr modified 7075 aluminum alloy prepared by selective laser melting Jiang Bi, Zhenglong Lei *, Yanbin Chen **, Xi Chen, Ze Tian, Jingwei Liang, Xinrui Zhang, Xikun Qin State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Selective laser melting 7075 aluminum alloy Energy density Microstructure Mechanical properties
A novel Sc- and Zr- modified 7075 Al alloy with a low microalloying content (0.4 wt% Sc and 0.25 wt %Zr) was designed and the influence of process parameters on the microstructure and mechanical properties of selective laser melting (SLM) samples was systematically analyzed. As the SLM energy input increased, the relative density of block specimens first increased before plateauing. At a high energy density, crack defects disappeared, and the average grain size significantly decreased. For the specimen fabricated at 375 J/mm3, the average grain size was 2.6 μm, which is only 9.8% of the size of samples fabricated at 44 J/mm3. Due to fine grain strengthening, the mechanical properties of the printed specimens were remarkably improved, but the high energy input softened the matrix. Due to these two opposing effects, the compressive strength and nano-hardness of specimens fabricated at 375 J/mm3 were 621 MPa and 1.85 GPa, which are respectively 129.6% and 98.4% of the specimen fabricated at 44 J/mm3.
1. Introduction New processing technologies are constantly emerging to meet the requirements of modern industry. In order to realize the integration of structure and function of parts, complex structures and personalized customization have become the preferred method of parts processing [1]. Additive manufacturing (AM) makes it easy and feasible to manu facture complex components as a whole [2–4]. SLM is a novel 3D printing technology that can be used to precisely manufacture metal parts (Fe, Ni, Al, Ti et al.) layer-by-layer [5–8]. Compared with other traditional processing methods (e.g., casting [9], stamping [10], extru sion [11] and rolling [12]), SLM has several advantages including fine microstructure, excellent mechanical properties, a short processing cycle and near net shape production [13,14]. Therefore, this technology has attracted significant in the literature and has gradually become an especially popular topic in AM research. Aluminum (Al) alloys are the most commonly used structural ma terials in various engineering fields, and are widely applied in aviation, aerospace, and automobile industries due to their excellent combination of high specific strength, low thermal expansion coefficients, and good corrosion resistance [15]. Thus, materials containing Al alloys with
complex structures are broadly used in industrial applications [16]. Furthermore, due to the high energy density and machining precision of SLM, Al parts produced using this method can be directly formed without secondary processing [17]. However, based on the special ma terial characteristics of Al alloys, such as low melting point, high thermal conductivity, and high laser reflectivity, the fabrication of Al parts using SLM is more difficult than other materials [18,19]. The SLM processing of Al alloys has been extensively investigated. For example, Zhang et al. [20] summarized the advantages and defects of the SLM processing of Al alloys, and suggested that the comprehensive properties of Al parts should be further investigated. Yang et al. [21] studied the effect of linear energy density on the vertical surface roughness of AlSi10Mg parts obtained using SLM, and concluded that the roughness decreased from 15 μm to 4 μm at a linear energy of 5.7 J/cm. Yang et al. [22] investigated the origin of porosity during the SLM process of Al–Si alloys and compared the influence of porosity location on the fatigue properties of SLM samples. The experimental results showed that porosity defects at the contour scanning region and at the edge of the core were detrimental to the fatigue properties of the materials. Wang et al. [23] studied the microstructure and mechanical properties of Al–Cu alloys prepared by SLM and found that the tensile
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J. Bi), [email protected] (Z. Lei), [email protected] (Y. Chen). https://doi.org/10.1016/j.msea.2019.138478 Received 11 August 2019; Accepted 29 September 2019 Available online 30 September 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Materials Science & Engineering A 768 (2019) 138478
Fig. 1. SEM image of Sc, Zr modified AA7075 powders (a) and the particle size range (b), AA7075 powder (c) and magnification of Zone D in 1c (d). Table 1 Chemical composition of Sc–Zr modified AA7075 powder in wt%. Elements
Si
Mg
Fe
Cu
Cr
Ti
Mn
Zn
Sc
Zr
Al
Actual
0.2
2.5
0.2
1.6
0.2
0.01
0.05
5.8
0.4
0.25
Bal
et al. [29–31] systematically investigated the influence of Zr addition on the microstructure and mechanical properties of a 2024 Al alloy. Zhou et al. [32] designed a new Sc- and Zr- modified Al–Zn–Mg alloys for SLM, which exhibited good mechanical properties after T6 heat treatment. Yang et al. [33] added 1.08 wt% Sc to an Al–Mg–Zr alloy and found that the grain size was significantly refined. Currently, research on the SLM of Sc- and Zr- modified Al alloys has mainly concentrated on 5xxx series alloys [34–38], and only AlMgScZr (5024) powder has been commercialized. However, high strength Al powders (2xxx, 6xxx and 7xxx series Al alloys) with different Sc/Zr ra tios are typically custom-designed at the lab-scale, and are consequently extremely expensive to manufacture. Furthermore, Sc is very expensive and the optimum Sc content for powders is still unknown, but the most economical method is to add microalloying elements by mechanical mixing [6,24,39]. The aim of this paper is to design a new Sc- and Zrmodified 7075 Al alloy with low Sc and Zr contents using ball milling and systematically studying the influence of process parameters on the
Table 2 Process parameters of block specimens. Process parameter
Value
Laser power (W) Scanning speed (mm/s) Layer thickness (μm) Hatching space (μm)
200, 240, 280, 320, 360 400, 800, 1200, 1600, 2000 30 80, 100, 120
strength reached to 455 MPa after T6 (solution and aging) heat treat ment. By systematically studying of SLM processing of Al alloys, it was found that the 2xxx, 6xxx, and 7xxx series Al alloys were nearly impossible to print due to thermal cracking [24–27]. However, when powders were micro-alloyed with Sc and Zr, many heterogeneous nucleation particles were formed during printing, and crack defects were effectively eliminated due to grain refinement. As a result, the mechanical properties of the printed parts were improved [28]. Zhu
Fig. 2. Block specimens (a) and scanning strategy (b). 2
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spherical powder particles, the weight ratio of ball-to-powder was set as 1:1. According to previous studies, the weight ratio of AA7075, pure Sc, and ZrAl16 powders was set as 99.35: 0.4: 0.25. The AA7075 powder (d10 ¼ 23.46 μm, d50 ¼ 41.37 μm, d90 ¼ 58.58 μm) (Fig. 1(b)) was pro duced by TLS Technik, and the microalloying powders (ZrAl16 (~37 μm) and pure Sc (~74 μm)) were produced by Jinzhou Haixin Metal Materials Co., Ltd, respectively. After ball milling process, the composite powders were evenly mixed (Fig. 1(a)). Some planetary powders were distributed on the surface of the AA7075 powder (Fig. 1 (c)). In the magnified zone D in Fig. 1(c), the grain size of the AA7075 powder was approximately 2–5 μm (Fig. 1(d)). The chemical composi tion of the mixed powders is shown in Table 1. 2.2. SLM process The experimental process was performed using an EOS M290 equipped with a Yb fiber laser (wavelength: 1064 nm; laser power: 0–400 W; spot diameter: 100 μm) and an option scanner system (scan ning speed: 0–7000 mm/s). A total of 75 block specimen were prepared to determine the effect of process parameters on the density of SLM samples (Fig. 2). The process parameters were listed in Table 2. During printing, the oxygen content in the forming chamber was set to less than 0.1%, and the energy density (E) of the SLM samples was calculated using Eq. (1).
Fig. 3. Relative density of block specimens at different energy densities.
microstructure and mechanical behavior of SLM samples. The Sc- and Zr- modified 7xxx series Al alloys will further promote the application of lightweight, complex parts in aerospace engineering. 2. Experimental material and methods
E ¼ P/vht
2.1. Materials
where P is the laser power, v is the scanning speed, h is the scanning hatch and t is the layer thickness.
The experimental material used for SLM was Sc- and Zr- modified 7075 Al alloy powder - prepared by a -MITR-YXQM-2L planetary ball mill (rotation speed: 200 rpm; mixing time: 4 h) [40]. To ensure
2.3. Phase identification and morphological characterization
(1)
Microstructure observations were performed using an optical
Fig. 4. DSC curves (a) and XRD patterns (b) of powder and block specimens.
Fig. 5. Microstructure of the SLM sample: (a) XY plane and (b) YZ plane. 3
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Fig. 6. EBSD images of block specimens at different energy densities: (a) 44 J/mm3, (b) 88 J/mm3, (c) 167 J/mm3, (d) 222 J/mm3, (e) 375 J/mm3 and (f) magnified zone F in (e).
Fig. 7. EPMA analysis of the block specimens prepared at different energy densities: (a)–(d) 88 J/mm3, (e)–(h) 375 J/mm3.
microscope (OM, VHX-1000, KEYENCE, Japan), a scanning electron microscope (SEM, Quanta 200FEG, FEI, USA) and a transmission elec tron microscope (TEM, Talos F200x, FEI, USA). OM samples were pol ished and etched by Keller’s reagent for 10 s [41]. The grain size and
crystal orientation of samples were further investigated by electron back scattering diffraction (EBSD). EBSD samples were prepared by electro polishing in an electrolyte (ethanol and perchloric acid in a 9:1 ratio) at 20 � C [10]. The phases of powders and block specimens were analyzed 4
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properties of block specimens, their nanohardess and compressive properties were systematically studied. The compressive properties of SLM samples were evaluated by an electronic universal testing machine (AGX-plus, Japan) at room temperature. Nanohardness was measured by a nanoindentation tester (G200) at a constant load of 8000 μN and a hold time of 2 s. Ten hardness readings were recorded per specimen and the distance between two adjacent test points was set at 20 μm. 3. Results and discussion 3.1. Densification Fig. 3 illustrates the relative densities of block specimens under different energy densities (27.8–375 J/mm3). At an energy density of 27.8 J/mm3, the density was low and the porosity of the samples reached 18.2% due to insufficient powder melting. As a result, many unfused voids formed inside the SLM sample. In the energy density range from 27.8 to 80 J/mm3, the relative densities of the block specimens increased significantly with the energy density. The maximum relative density was 99.2% at an energy input of 72.9 J/mm3, and as the energy density further increased, the relative densities of the block specimens decreased slightly. At high energy densities, the temperature of the melt pool was notably higher than samples obtained at a low energy density. Elements with low melting points easily burned and produced the metal vapors. Moreover, the SLM scanning speed was so fast that the cooling rate was about 106–108 K/s [5]. Due to rapid solidification, the metal vapor barely escaped the melt pool, and formed pore defects during SLM. 3.2. XRD and DSC analyses Fig. 4 shows the differential scanning calorimetry (DSC) curves and XRD patterns of powder and block specimens. Fig. 4(a) shows the DSC trace of AA7075 powder and an SLM sample obtained from mixed powders. The temperature range was 25–700 � C with a heating rate of 10 � C/min. From the DSC curves, it can be seen that the melting point of the AA7075 powder and block samples were 644.5 and 656.3 � C, respectively. In Fig. 4(b), α-Al was detected in the powder and block samples. During SLM, the mixed powders melted and then solidified. Due to the microalloying, Sc and Zr reacted with the Al matrix, which formed an Al3(Sc, Zr) phase that was coherent with the Al matrix. 3.3. Microstructure observation 3.3.1. Etched microstructure Fig. 5 shows the typical microstructure of an SLM sample fabricated at 111 J/mm3. At the top surface (XY plane) of the specimen, some vertical cross laser scanning traces can be observed (Fig. 5(a)). The laser scanning trace was discontinuous and had a width ranging from 70 to 160 μm, with some fine grains distributed at its boundary. In Fig. 5(b), some pores and microcracks were present on the YZ plane of the block specimen. The morphology of the melt pool was clearly visible and showed a regular scaly shape. In the center of the melt pool, the microstructure was comprised of coarse columnar grains. However, at the melt pool boundary, some fine equiaxed grains formed because the temperature was slightly higher that melting point of the Al alloy (far below the temperature at the center of the molten pool) [43]. Due to the combined effects of fluid flow and heat transfer, the Al3(Sc, Zr) particles were formed, which effectively served as heterogeneous nucleation sites, and obvious grain refinement was observed at the melt pool boundary.
Fig. 8. The element distribution of SLM sample: (a)–(b) HADDF image of the SLM sample; (c)–(h) distribution analysis of Al, Zn, Mg, Cu, Sc, and Zr.
by X-ray diffraction (XRD, Empyrean, Panalytical, Netherlands) with diffraction angles ranging from 20� to 100� at a scan rate of 5� /min and a step size of 0.02� [42]. The element distribution was measured by electron probe microanalysis (EPMA, JXA-8230, Japan) and high angular annular dark field (HAADF) imaging.yTEM specimens were first ground to 50 μm and then thinned using an ion beam thinner (GATAN695) until they were penetrated.
3.3.2. Grain refinement Fig. 6 shows the orientation maps of SLM samples at different energy densities. In Fig. 6(a), the sample has a coarse columnar crystal micro structure, with an average grain size of 26.4 μm (44 J/mm3). As the
2.4. Mechanical properties To investigate the influence of Sc and Zr contents on the mechanical 5
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Fig. 9. TEM analysis of the SLM samples fabricated at an energy density of 375 J/mm3: (a) bright field image (b) precipitates inside the grain; (c) composite precipitates; and (d) HRTEM of zone D.
Fig. 10. Precipitate in the Al matrix (a) and the line scanning result (b).
energy density increased, the grain refinement became increasingly obvious. In the sample prepared at 88 J/mm3, the columnar crystal grain size decreased to 8.1 μm (Fig. 6(b)). In Fig. 6(c), it can be seen that some columnar crystals transformed into equiaxed crystals with fine equiaxed crystals mainly distributed at the melt pool boundary. As the energy density further increased, all columnar crystals disappeared and trans formed into equiaxed dendrites (Fig. 6(d)). As for the sample fabricated at an energy density of 375 J/mm3, the microstructure primarily con tained equiaxed crystals with an average grain size of 2.6 μm. The mi crostructures were divided into two distinct regions - the coarse grain
region and fine grain region (Fig. 6(e)) - which is similar to the result of Ref. [33]. In the magnified zone F, submicron grains were observed ((Fig. 6(f)). 3.3.3. Element distribution From Fig. 6, significant grain refinement was observed at high en ergy densities. In order to obtain determine the reason for this phe nomenon, the elemental distribution of the block sample was tested by EMPA, as shown in Fig. 7. Fig. 7(a)–(d) contains the BSE image and elemental distribution of Al, Sc, and Zr at a low energy density (88 J/ 6
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Fig. 11. Precipitate and line scanning result (a), the corresponding HRTEM image of zone B (b), zone C (c) and zone D (d).
Fig. 12. Loading and unloading curves (a); nanohardness and modulus (b).
mm3), and Fig. 7(e)–(h) are the corresponding high energy density (375 J/mm3). In Fig. 7(a), the contour of the melt pool is clear visible, with a width of about 182 μm. In the sample exposed to 375 J/mm3, the boundary of the melt pool is blurred, and the microstructure contains fine equiaxed grains (Fig. 7(e)). As for the element distribution, Sc and Zr aggregated in the Al matrix (Fig. 7(b)–(d)), while the microalloying elements were evenly distributed as the energy density increased to 375 J/mm3 (Fig. 7(f)–(h)). From the Al–Sc and Al–Zr binary alloy phase diagrams, the melting points of Sc and ZrAl16 were 1541 and 1419 � C. During SLM, the temperature of the melt pool, which was between melting point and boiling point of Al alloy, increased with the energy
density. At an energy density of 88 J/mm3, the temperature of the melt pool was lower than the melting point of Sc and ZrAl16, preventing them from dissolving in the melt pool. The temperature increased signifi cantly with the energy input, and exceeded the melting point of the Sc and ZrAl16 powders. Sc and Zr reacted with Al and the Al3(Sc, Zr) phase generated during the solidification process, which dispersed Sc and Zr throughout the Al matrix. Fig. 8 shows the high angular annular dark field (HADDF) images of SLM samples fabricated at the energy density of 375 J/mm3. In Fig. 8(a), many eutectic precipitated phases were located at the grain boundary and some fine spherical precipitates (100–200 nm) were distributed 7
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Fig. 13. Compressive stress-strain curves (a) 400 mm/s and (b) compressive strength and strain.
inside the grains. The size of precipitated phase at the grain boundary ranged from 0.6 to 1.81 μm. The element distributions in Fig. 8(c)–(f) show that Al was mainly distributed in the inner grains (Fig. 8(c)), while Zn, Mg, and Cu were distributed at the grain boundary. After adding Sc and Zr elements, the Al3(Sc, Zr) particles formed which acted as het erogeneous nucleation sites (Fig. 8(g) and (h)), and the grain size decreased significantly due to grain refinement.
Fig. 12(b) shows the modulus and nanohardess of the SLM samples. During testing, the maximum load was kept constant at 8 μN. Since the loading was constant, a higher nanohardess resulted in a smaller indentation depth. In Fig. 12(b), the nanohardness is 1.88 GPa and 1.85 GPa at energy densities of 44 J/mm3 and 375 J/mm3, respectively. The largest nanohardness was 2.04 GPa (222 J/mm3), which is 10.3% higher than the sample prepared at an energy density is 375 J/mm3. At high energy densities, the average grain size was smaller at low energy densities. Despite grain refinement, the high energy input caused ma terial softening, and these two combined effects caused the nanohard ness to first increase and then decreases with increasing energy density.
3.3.4. Phase analysis Fig. 9 shows the TEM analysis of the samples fabricated at an SLM energy density of 375 J/mm3. Some fine sub-grains with grain sizes of 200–500 nm were formed during SLM (Fig. 9(a)). The bright field image shows that elongated eutectic precipitates were presented at grain boundaries, and some spherical precipitates formed inside grains. Despite the presence of these precipitates, many dislocations were entangled in the inner grains (Fig. 9(b)). In Fig. 9(c), different precipi tate phases aggregated to form a composite precipitate phase. The highresolution TEM (HRTEM) image in Fig. 9(d) further shows that the composite precipitate phase in zone D was formed by multiple precipitates. To further investigate the composition of the precipitated phase in Fig. 9(c), energy dispersive spectroscopy (EDS) was performed on the composite precipitates, and the EDS line result is shown in Fig. 10. In Figs. 9(c) and 10(a), the positions of different phases can be distin guished by the different colors of precipitates. From the ESD line result, the element content of the measured line was obtained, allowing the types of precipitates to be determined. The main elements of 7075 Al alloy were Zn, Mg, and Cu, while common precipitates in the Al matrix were η-MgZn2, and θ-Al2Cu. After adding Sc and Zr microalloying ele ments, the Al3(Sc, Zr) phase precipitated and dispersed throughout the Al matrix during solidification. The mechanical properties were signif icantly improved due to the dispersed distribution of precipitates in the matrix. Fig. 11 shows the cross section of a precipitate in the inner grain. Fig. 11(a) shows the EDS line scanning result of the composite phase and Fig. 11(b) shows the HRTEM of zone B s in Fig. 11(a). By analyzing of Fig. 11(a) and (b), the precipitate was determined to be a composite phase. According to the line scanning result, the main elements in zone C were Al and Cu, while zone D primarily contained Al, Sc, and Zr. Combined with the HRTEM images of zone C and zone D (Fig. 11(c) and (d)), the precipitated phases were θ-Al2Cu and Al3(Sc, Zr). Al2Cu had a relatively complicated crystal structure (a ¼ 0.6063 nm, c ¼ 0.4872 nm, I4/mcm (no. 140) space group and tI12 symmetry) [44]. The Al3(Sc, Zr) phase is coherent with the Al matrix.
3.4.2. Compressive properties Fig. 13 shows the compressive properties of the block samples under different energy densities. Fig. 13(a) shows the compressive curves ob tained at a strain rate of 1 mm/min. The strain hardening exponents of the different specimens were nearly identical, but the compress strain increased with increasing energy density. As previously mentioned, grain refinement was significant at high energy densities, and the compressive strain gradually increased as the grain size decreased. In Fig. 13(b), the compressive strength increased as the energy density increased. At an energy of 44 J/mm3, the compressive strength was 479 MPa, and when the energy density was 375 J/mm3, the compressive strength reached its maximum value of 621 MPa, which was 29.6% higher than that of 44 J/mm3. 4. Conclusions In this paper, a novel Sc- and Zr- modified 7075 Al alloy was designed for industrial applications. A total of 75 block specimens were fabricated at different energy densities and the relative densities of these samples were obtained. The influence of process parameters on the microstructure and mechanical properties of block specimens was sys tematically studied, and the following conclusions were drawn: (1) The relative density of block specimens first increased and then plateaued as the energy density increased. Very dense samples were prepared within the energy density range of 68–375 J/mm3. The maximum relative density (99.2%) was obtained at an energy density of 72.9 J/mm3. (2) Energy density significantly affected the grain morphology of samples obtained using SLM. At a low energy density (44 J/ mm3), a coarse columnar crystal microstructure was obtained, with an average grain size of 26.4 μm. The microstructure of the sample obtained at an energy density of 375 J/mm3 contained equiaxed crystals, and the average grain size decreased to 2.6 μm, which was 9.8% of the specimen fabricated at 44 J/mm3. (3) At a low energy input (88 J/mm3), the temperature of the melt pool was so low that the Sc and ZrAl16 powders could not dissolve and sufficiently diffuse throughout the Al matrix, while
3.4. Mechanical properties 3.4.1. Nanohardness Fig. 12 shows nanohardness of block samples prepared at different energy densities. Fig. 12(a) shows the loading and unloading curves, and 8
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the Sc and Zr were evenly distributed (at an energy density of 375 J/mm3) due to fluid flow and heat transfer. As an effective heterogeneous nucleating agent, the Al3(Sc, Zr) phase precipi tated during solidification due to a reaction with the Al matrix. As a result, the grain size of the SLM specimen was obviously refined. (4) With the combined effect of fine grain strengthening and matrix softening at high energy densities, the compressive strength and nano-hardness of specimen fabricated at 375 J/mm3 are 621 MPa and 1.85 GPa, which are 129.6% and 98.4% of the specimen fabricated at 44 J/mm3.
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Microstructure and mechanical properties of a novel Sc and Zr modified 7075 aluminum alloy prepared by selective laser melting Jiang Bi, Zhenglong Lei *, Yanbin Chen **, Xi Chen, Ze Tian, Jingwei Liang, Xinrui Zhang, Xikun Qin State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Selective laser melting 7075 aluminum alloy Energy density Microstructure Mechanical properties
A novel Sc- and Zr- modified 7075 Al alloy with a low microalloying content (0.4 wt% Sc and 0.25 wt %Zr) was designed and the influence of process parameters on the microstructure and mechanical properties of selective laser melting (SLM) samples was systematically analyzed. As the SLM energy input increased, the relative density of block specimens first increased before plateauing. At a high energy density, crack defects disappeared, and the average grain size significantly decreased. For the specimen fabricated at 375 J/mm3, the average grain size was 2.6 μm, which is only 9.8% of the size of samples fabricated at 44 J/mm3. Due to fine grain strengthening, the mechanical properties of the printed specimens were remarkably improved, but the high energy input softened the matrix. Due to these two opposing effects, the compressive strength and nano-hardness of specimens fabricated at 375 J/mm3 were 621 MPa and 1.85 GPa, which are respectively 129.6% and 98.4% of the specimen fabricated at 44 J/mm3.
1. Introduction New processing technologies are constantly emerging to meet the requirements of modern industry. In order to realize the integration of structure and function of parts, complex structures and personalized customization have become the preferred method of parts processing [1]. Additive manufacturing (AM) makes it easy and feasible to manu facture complex components as a whole [2–4]. SLM is a novel 3D printing technology that can be used to precisely manufacture metal parts (Fe, Ni, Al, Ti et al.) layer-by-layer [5–8]. Compared with other traditional processing methods (e.g., casting [9], stamping [10], extru sion [11] and rolling [12]), SLM has several advantages including fine microstructure, excellent mechanical properties, a short processing cycle and near net shape production [13,14]. Therefore, this technology has attracted significant in the literature and has gradually become an especially popular topic in AM research. Aluminum (Al) alloys are the most commonly used structural ma terials in various engineering fields, and are widely applied in aviation, aerospace, and automobile industries due to their excellent combination of high specific strength, low thermal expansion coefficients, and good corrosion resistance [15]. Thus, materials containing Al alloys with
complex structures are broadly used in industrial applications [16]. Furthermore, due to the high energy density and machining precision of SLM, Al parts produced using this method can be directly formed without secondary processing [17]. However, based on the special ma terial characteristics of Al alloys, such as low melting point, high thermal conductivity, and high laser reflectivity, the fabrication of Al parts using SLM is more difficult than other materials [18,19]. The SLM processing of Al alloys has been extensively investigated. For example, Zhang et al. [20] summarized the advantages and defects of the SLM processing of Al alloys, and suggested that the comprehensive properties of Al parts should be further investigated. Yang et al. [21] studied the effect of linear energy density on the vertical surface roughness of AlSi10Mg parts obtained using SLM, and concluded that the roughness decreased from 15 μm to 4 μm at a linear energy of 5.7 J/cm. Yang et al. [22] investigated the origin of porosity during the SLM process of Al–Si alloys and compared the influence of porosity location on the fatigue properties of SLM samples. The experimental results showed that porosity defects at the contour scanning region and at the edge of the core were detrimental to the fatigue properties of the materials. Wang et al. [23] studied the microstructure and mechanical properties of Al–Cu alloys prepared by SLM and found that the tensile
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J. Bi), [email protected] (Z. Lei), [email protected] (Y. Chen). https://doi.org/10.1016/j.msea.2019.138478 Received 11 August 2019; Accepted 29 September 2019 Available online 30 September 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Fig. 1. SEM image of Sc, Zr modified AA7075 powders (a) and the particle size range (b), AA7075 powder (c) and magnification of Zone D in 1c (d). Table 1 Chemical composition of Sc–Zr modified AA7075 powder in wt%. Elements
Si
Mg
Fe
Cu
Cr
Ti
Mn
Zn
Sc
Zr
Al
Actual
0.2
2.5
0.2
1.6
0.2
0.01
0.05
5.8
0.4
0.25
Bal
et al. [29–31] systematically investigated the influence of Zr addition on the microstructure and mechanical properties of a 2024 Al alloy. Zhou et al. [32] designed a new Sc- and Zr- modified Al–Zn–Mg alloys for SLM, which exhibited good mechanical properties after T6 heat treatment. Yang et al. [33] added 1.08 wt% Sc to an Al–Mg–Zr alloy and found that the grain size was significantly refined. Currently, research on the SLM of Sc- and Zr- modified Al alloys has mainly concentrated on 5xxx series alloys [34–38], and only AlMgScZr (5024) powder has been commercialized. However, high strength Al powders (2xxx, 6xxx and 7xxx series Al alloys) with different Sc/Zr ra tios are typically custom-designed at the lab-scale, and are consequently extremely expensive to manufacture. Furthermore, Sc is very expensive and the optimum Sc content for powders is still unknown, but the most economical method is to add microalloying elements by mechanical mixing [6,24,39]. The aim of this paper is to design a new Sc- and Zrmodified 7075 Al alloy with low Sc and Zr contents using ball milling and systematically studying the influence of process parameters on the
Table 2 Process parameters of block specimens. Process parameter
Value
Laser power (W) Scanning speed (mm/s) Layer thickness (μm) Hatching space (μm)
200, 240, 280, 320, 360 400, 800, 1200, 1600, 2000 30 80, 100, 120
strength reached to 455 MPa after T6 (solution and aging) heat treat ment. By systematically studying of SLM processing of Al alloys, it was found that the 2xxx, 6xxx, and 7xxx series Al alloys were nearly impossible to print due to thermal cracking [24–27]. However, when powders were micro-alloyed with Sc and Zr, many heterogeneous nucleation particles were formed during printing, and crack defects were effectively eliminated due to grain refinement. As a result, the mechanical properties of the printed parts were improved [28]. Zhu
Fig. 2. Block specimens (a) and scanning strategy (b). 2
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spherical powder particles, the weight ratio of ball-to-powder was set as 1:1. According to previous studies, the weight ratio of AA7075, pure Sc, and ZrAl16 powders was set as 99.35: 0.4: 0.25. The AA7075 powder (d10 ¼ 23.46 μm, d50 ¼ 41.37 μm, d90 ¼ 58.58 μm) (Fig. 1(b)) was pro duced by TLS Technik, and the microalloying powders (ZrAl16 (~37 μm) and pure Sc (~74 μm)) were produced by Jinzhou Haixin Metal Materials Co., Ltd, respectively. After ball milling process, the composite powders were evenly mixed (Fig. 1(a)). Some planetary powders were distributed on the surface of the AA7075 powder (Fig. 1 (c)). In the magnified zone D in Fig. 1(c), the grain size of the AA7075 powder was approximately 2–5 μm (Fig. 1(d)). The chemical composi tion of the mixed powders is shown in Table 1. 2.2. SLM process The experimental process was performed using an EOS M290 equipped with a Yb fiber laser (wavelength: 1064 nm; laser power: 0–400 W; spot diameter: 100 μm) and an option scanner system (scan ning speed: 0–7000 mm/s). A total of 75 block specimen were prepared to determine the effect of process parameters on the density of SLM samples (Fig. 2). The process parameters were listed in Table 2. During printing, the oxygen content in the forming chamber was set to less than 0.1%, and the energy density (E) of the SLM samples was calculated using Eq. (1).
Fig. 3. Relative density of block specimens at different energy densities.
microstructure and mechanical behavior of SLM samples. The Sc- and Zr- modified 7xxx series Al alloys will further promote the application of lightweight, complex parts in aerospace engineering. 2. Experimental material and methods
E ¼ P/vht
2.1. Materials
where P is the laser power, v is the scanning speed, h is the scanning hatch and t is the layer thickness.
The experimental material used for SLM was Sc- and Zr- modified 7075 Al alloy powder - prepared by a -MITR-YXQM-2L planetary ball mill (rotation speed: 200 rpm; mixing time: 4 h) [40]. To ensure
2.3. Phase identification and morphological characterization
(1)
Microstructure observations were performed using an optical
Fig. 4. DSC curves (a) and XRD patterns (b) of powder and block specimens.
Fig. 5. Microstructure of the SLM sample: (a) XY plane and (b) YZ plane. 3
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Fig. 6. EBSD images of block specimens at different energy densities: (a) 44 J/mm3, (b) 88 J/mm3, (c) 167 J/mm3, (d) 222 J/mm3, (e) 375 J/mm3 and (f) magnified zone F in (e).
Fig. 7. EPMA analysis of the block specimens prepared at different energy densities: (a)–(d) 88 J/mm3, (e)–(h) 375 J/mm3.
microscope (OM, VHX-1000, KEYENCE, Japan), a scanning electron microscope (SEM, Quanta 200FEG, FEI, USA) and a transmission elec tron microscope (TEM, Talos F200x, FEI, USA). OM samples were pol ished and etched by Keller’s reagent for 10 s [41]. The grain size and
crystal orientation of samples were further investigated by electron back scattering diffraction (EBSD). EBSD samples were prepared by electro polishing in an electrolyte (ethanol and perchloric acid in a 9:1 ratio) at 20 � C [10]. The phases of powders and block specimens were analyzed 4
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properties of block specimens, their nanohardess and compressive properties were systematically studied. The compressive properties of SLM samples were evaluated by an electronic universal testing machine (AGX-plus, Japan) at room temperature. Nanohardness was measured by a nanoindentation tester (G200) at a constant load of 8000 μN and a hold time of 2 s. Ten hardness readings were recorded per specimen and the distance between two adjacent test points was set at 20 μm. 3. Results and discussion 3.1. Densification Fig. 3 illustrates the relative densities of block specimens under different energy densities (27.8–375 J/mm3). At an energy density of 27.8 J/mm3, the density was low and the porosity of the samples reached 18.2% due to insufficient powder melting. As a result, many unfused voids formed inside the SLM sample. In the energy density range from 27.8 to 80 J/mm3, the relative densities of the block specimens increased significantly with the energy density. The maximum relative density was 99.2% at an energy input of 72.9 J/mm3, and as the energy density further increased, the relative densities of the block specimens decreased slightly. At high energy densities, the temperature of the melt pool was notably higher than samples obtained at a low energy density. Elements with low melting points easily burned and produced the metal vapors. Moreover, the SLM scanning speed was so fast that the cooling rate was about 106–108 K/s [5]. Due to rapid solidification, the metal vapor barely escaped the melt pool, and formed pore defects during SLM. 3.2. XRD and DSC analyses Fig. 4 shows the differential scanning calorimetry (DSC) curves and XRD patterns of powder and block specimens. Fig. 4(a) shows the DSC trace of AA7075 powder and an SLM sample obtained from mixed powders. The temperature range was 25–700 � C with a heating rate of 10 � C/min. From the DSC curves, it can be seen that the melting point of the AA7075 powder and block samples were 644.5 and 656.3 � C, respectively. In Fig. 4(b), α-Al was detected in the powder and block samples. During SLM, the mixed powders melted and then solidified. Due to the microalloying, Sc and Zr reacted with the Al matrix, which formed an Al3(Sc, Zr) phase that was coherent with the Al matrix. 3.3. Microstructure observation 3.3.1. Etched microstructure Fig. 5 shows the typical microstructure of an SLM sample fabricated at 111 J/mm3. At the top surface (XY plane) of the specimen, some vertical cross laser scanning traces can be observed (Fig. 5(a)). The laser scanning trace was discontinuous and had a width ranging from 70 to 160 μm, with some fine grains distributed at its boundary. In Fig. 5(b), some pores and microcracks were present on the YZ plane of the block specimen. The morphology of the melt pool was clearly visible and showed a regular scaly shape. In the center of the melt pool, the microstructure was comprised of coarse columnar grains. However, at the melt pool boundary, some fine equiaxed grains formed because the temperature was slightly higher that melting point of the Al alloy (far below the temperature at the center of the molten pool) [43]. Due to the combined effects of fluid flow and heat transfer, the Al3(Sc, Zr) particles were formed, which effectively served as heterogeneous nucleation sites, and obvious grain refinement was observed at the melt pool boundary.
Fig. 8. The element distribution of SLM sample: (a)–(b) HADDF image of the SLM sample; (c)–(h) distribution analysis of Al, Zn, Mg, Cu, Sc, and Zr.
by X-ray diffraction (XRD, Empyrean, Panalytical, Netherlands) with diffraction angles ranging from 20� to 100� at a scan rate of 5� /min and a step size of 0.02� [42]. The element distribution was measured by electron probe microanalysis (EPMA, JXA-8230, Japan) and high angular annular dark field (HAADF) imaging.yTEM specimens were first ground to 50 μm and then thinned using an ion beam thinner (GATAN695) until they were penetrated.
3.3.2. Grain refinement Fig. 6 shows the orientation maps of SLM samples at different energy densities. In Fig. 6(a), the sample has a coarse columnar crystal micro structure, with an average grain size of 26.4 μm (44 J/mm3). As the
2.4. Mechanical properties To investigate the influence of Sc and Zr contents on the mechanical 5
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Fig. 9. TEM analysis of the SLM samples fabricated at an energy density of 375 J/mm3: (a) bright field image (b) precipitates inside the grain; (c) composite precipitates; and (d) HRTEM of zone D.
Fig. 10. Precipitate in the Al matrix (a) and the line scanning result (b).
energy density increased, the grain refinement became increasingly obvious. In the sample prepared at 88 J/mm3, the columnar crystal grain size decreased to 8.1 μm (Fig. 6(b)). In Fig. 6(c), it can be seen that some columnar crystals transformed into equiaxed crystals with fine equiaxed crystals mainly distributed at the melt pool boundary. As the energy density further increased, all columnar crystals disappeared and trans formed into equiaxed dendrites (Fig. 6(d)). As for the sample fabricated at an energy density of 375 J/mm3, the microstructure primarily con tained equiaxed crystals with an average grain size of 2.6 μm. The mi crostructures were divided into two distinct regions - the coarse grain
region and fine grain region (Fig. 6(e)) - which is similar to the result of Ref. [33]. In the magnified zone F, submicron grains were observed ((Fig. 6(f)). 3.3.3. Element distribution From Fig. 6, significant grain refinement was observed at high en ergy densities. In order to obtain determine the reason for this phe nomenon, the elemental distribution of the block sample was tested by EMPA, as shown in Fig. 7. Fig. 7(a)–(d) contains the BSE image and elemental distribution of Al, Sc, and Zr at a low energy density (88 J/ 6
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Fig. 11. Precipitate and line scanning result (a), the corresponding HRTEM image of zone B (b), zone C (c) and zone D (d).
Fig. 12. Loading and unloading curves (a); nanohardness and modulus (b).
mm3), and Fig. 7(e)–(h) are the corresponding high energy density (375 J/mm3). In Fig. 7(a), the contour of the melt pool is clear visible, with a width of about 182 μm. In the sample exposed to 375 J/mm3, the boundary of the melt pool is blurred, and the microstructure contains fine equiaxed grains (Fig. 7(e)). As for the element distribution, Sc and Zr aggregated in the Al matrix (Fig. 7(b)–(d)), while the microalloying elements were evenly distributed as the energy density increased to 375 J/mm3 (Fig. 7(f)–(h)). From the Al–Sc and Al–Zr binary alloy phase diagrams, the melting points of Sc and ZrAl16 were 1541 and 1419 � C. During SLM, the temperature of the melt pool, which was between melting point and boiling point of Al alloy, increased with the energy
density. At an energy density of 88 J/mm3, the temperature of the melt pool was lower than the melting point of Sc and ZrAl16, preventing them from dissolving in the melt pool. The temperature increased signifi cantly with the energy input, and exceeded the melting point of the Sc and ZrAl16 powders. Sc and Zr reacted with Al and the Al3(Sc, Zr) phase generated during the solidification process, which dispersed Sc and Zr throughout the Al matrix. Fig. 8 shows the high angular annular dark field (HADDF) images of SLM samples fabricated at the energy density of 375 J/mm3. In Fig. 8(a), many eutectic precipitated phases were located at the grain boundary and some fine spherical precipitates (100–200 nm) were distributed 7
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Fig. 13. Compressive stress-strain curves (a) 400 mm/s and (b) compressive strength and strain.
inside the grains. The size of precipitated phase at the grain boundary ranged from 0.6 to 1.81 μm. The element distributions in Fig. 8(c)–(f) show that Al was mainly distributed in the inner grains (Fig. 8(c)), while Zn, Mg, and Cu were distributed at the grain boundary. After adding Sc and Zr elements, the Al3(Sc, Zr) particles formed which acted as het erogeneous nucleation sites (Fig. 8(g) and (h)), and the grain size decreased significantly due to grain refinement.
Fig. 12(b) shows the modulus and nanohardess of the SLM samples. During testing, the maximum load was kept constant at 8 μN. Since the loading was constant, a higher nanohardess resulted in a smaller indentation depth. In Fig. 12(b), the nanohardness is 1.88 GPa and 1.85 GPa at energy densities of 44 J/mm3 and 375 J/mm3, respectively. The largest nanohardness was 2.04 GPa (222 J/mm3), which is 10.3% higher than the sample prepared at an energy density is 375 J/mm3. At high energy densities, the average grain size was smaller at low energy densities. Despite grain refinement, the high energy input caused ma terial softening, and these two combined effects caused the nanohard ness to first increase and then decreases with increasing energy density.
3.3.4. Phase analysis Fig. 9 shows the TEM analysis of the samples fabricated at an SLM energy density of 375 J/mm3. Some fine sub-grains with grain sizes of 200–500 nm were formed during SLM (Fig. 9(a)). The bright field image shows that elongated eutectic precipitates were presented at grain boundaries, and some spherical precipitates formed inside grains. Despite the presence of these precipitates, many dislocations were entangled in the inner grains (Fig. 9(b)). In Fig. 9(c), different precipi tate phases aggregated to form a composite precipitate phase. The highresolution TEM (HRTEM) image in Fig. 9(d) further shows that the composite precipitate phase in zone D was formed by multiple precipitates. To further investigate the composition of the precipitated phase in Fig. 9(c), energy dispersive spectroscopy (EDS) was performed on the composite precipitates, and the EDS line result is shown in Fig. 10. In Figs. 9(c) and 10(a), the positions of different phases can be distin guished by the different colors of precipitates. From the ESD line result, the element content of the measured line was obtained, allowing the types of precipitates to be determined. The main elements of 7075 Al alloy were Zn, Mg, and Cu, while common precipitates in the Al matrix were η-MgZn2, and θ-Al2Cu. After adding Sc and Zr microalloying ele ments, the Al3(Sc, Zr) phase precipitated and dispersed throughout the Al matrix during solidification. The mechanical properties were signif icantly improved due to the dispersed distribution of precipitates in the matrix. Fig. 11 shows the cross section of a precipitate in the inner grain. Fig. 11(a) shows the EDS line scanning result of the composite phase and Fig. 11(b) shows the HRTEM of zone B s in Fig. 11(a). By analyzing of Fig. 11(a) and (b), the precipitate was determined to be a composite phase. According to the line scanning result, the main elements in zone C were Al and Cu, while zone D primarily contained Al, Sc, and Zr. Combined with the HRTEM images of zone C and zone D (Fig. 11(c) and (d)), the precipitated phases were θ-Al2Cu and Al3(Sc, Zr). Al2Cu had a relatively complicated crystal structure (a ¼ 0.6063 nm, c ¼ 0.4872 nm, I4/mcm (no. 140) space group and tI12 symmetry) [44]. The Al3(Sc, Zr) phase is coherent with the Al matrix.
3.4.2. Compressive properties Fig. 13 shows the compressive properties of the block samples under different energy densities. Fig. 13(a) shows the compressive curves ob tained at a strain rate of 1 mm/min. The strain hardening exponents of the different specimens were nearly identical, but the compress strain increased with increasing energy density. As previously mentioned, grain refinement was significant at high energy densities, and the compressive strain gradually increased as the grain size decreased. In Fig. 13(b), the compressive strength increased as the energy density increased. At an energy of 44 J/mm3, the compressive strength was 479 MPa, and when the energy density was 375 J/mm3, the compressive strength reached its maximum value of 621 MPa, which was 29.6% higher than that of 44 J/mm3. 4. Conclusions In this paper, a novel Sc- and Zr- modified 7075 Al alloy was designed for industrial applications. A total of 75 block specimens were fabricated at different energy densities and the relative densities of these samples were obtained. The influence of process parameters on the microstructure and mechanical properties of block specimens was sys tematically studied, and the following conclusions were drawn: (1) The relative density of block specimens first increased and then plateaued as the energy density increased. Very dense samples were prepared within the energy density range of 68–375 J/mm3. The maximum relative density (99.2%) was obtained at an energy density of 72.9 J/mm3. (2) Energy density significantly affected the grain morphology of samples obtained using SLM. At a low energy density (44 J/ mm3), a coarse columnar crystal microstructure was obtained, with an average grain size of 26.4 μm. The microstructure of the sample obtained at an energy density of 375 J/mm3 contained equiaxed crystals, and the average grain size decreased to 2.6 μm, which was 9.8% of the specimen fabricated at 44 J/mm3. (3) At a low energy input (88 J/mm3), the temperature of the melt pool was so low that the Sc and ZrAl16 powders could not dissolve and sufficiently diffuse throughout the Al matrix, while
3.4. Mechanical properties 3.4.1. Nanohardness Fig. 12 shows nanohardness of block samples prepared at different energy densities. Fig. 12(a) shows the loading and unloading curves, and 8
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the Sc and Zr were evenly distributed (at an energy density of 375 J/mm3) due to fluid flow and heat transfer. As an effective heterogeneous nucleating agent, the Al3(Sc, Zr) phase precipi tated during solidification due to a reaction with the Al matrix. As a result, the grain size of the SLM specimen was obviously refined. (4) With the combined effect of fine grain strengthening and matrix softening at high energy densities, the compressive strength and nano-hardness of specimen fabricated at 375 J/mm3 are 621 MPa and 1.85 GPa, which are 129.6% and 98.4% of the specimen fabricated at 44 J/mm3.
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