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Microstructures and properties of an equimolar AlCoCrFeNi high entropy alloy printed by selective laser melting
Intermetallics 104 (2019) 24–32
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Microstructures and properties of an equimolar AlCoCrFeNi high entropy alloy printed by selective laser melting
T
P.D. Niua, R.D. Lia,b,∗, T.C. Yuana, S.Y. Zhuc, C. Chena, M.B. Wanga, L. Huanga,∗∗ a
State Key Laboratory of Powder Metallurgy, Science and Technology on High Strength Structural Materials Laboratory, Central South University, Changsha, 410083, PR China b State Key Laboratory of Solidification Processing Technology, Northwestern Polytechnical University, Xi'an, 710072, PR China c Materials Institute of Genome, Shanghai University, Shanghai, 200444, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: High entropy alloys Selective laser melting Equiatomic AlCoCrFeNi
The equimolar AlCoCrFeNi high entropy alloy (HEA) was printed by selective laser melting (SLM), with emphasis on its densification, phase identification, non-equilibrium microstructure and properties. The density of SLM samples increase gradually with the increase of volumetric energy density (VED) with the maximal relative density 98.4%. SLM samples are consisted of disordered body-center cubic phases (A2) and ordered body-center cubic phases (B2) which are different from the conventional cast or deformation with A2 and fcc phase; the content of B2 phase increases with the VED. Interestingly, the Fe-Cr precipitate, which has never been found in the cast one, is observed in the SLM one without heat treatment. The microstructure of SLM samples shows epitaxial growth of columnar A2 grain bundles perpendicular to the melt pool boundary with an average grain size of about 1.5 μm; the B2 phase is between the columnar A2 grains. At low VED, the SLM sample shows < 001 > preferred orientation, and is gradually transformed to the < 111 > crystallographic direction families at a higher VED. Besides, there exists the {100}A2∥{100}B2 orientation relationship in A2 and B2 phase. The maximal micro-hardness of 632.8Hv of SLM sample is higher than the cast one and the electron beam melting printed one.
1. Introduction The development of new materials is very important for the material science and the ecological utilization of limited resources on the earth. Conventional alloys like Al- alloy, Fe-alloy, Ti-alloys, TiAl-alloys, FeAlalloys are based on one or two main elements, and a relatively small amount of other elements are added to modify the microstructure and properties of the pure metal base. However, as a new class of metal alloy, high entropy alloys (HEAs) which provide new strategies and expand the range of traditional alloys for alloy design [1] attract widespread attention in recent years due to its excellent mechanical, high thermal stability, and outstanding wear resistance [2]. The concept of HEA was first proposed by Yeh et al. in 2004 [3]. In most of the cases, HEAs which are defined as alloys containing five or more major alloying elements possess a concentration between 5 at% and 35 at%, forming a single or double phase crystalline structure despite that there are multiple elements with different crystal structures. In particular, the quinary HEA composed of Al, Co, Cr, Fe, and Ni elements has received
∗
much attention due to its excellent properties such as good corrosion resistance, high hardness and high yield strength even at high temperatures [4]. Usually, HEA is manufactured by conventional casting or plastic deformation [5]; however, these methods fail to machine complex shapes and ultrafine grained structure. Therefore, for a HEA which has a promising application, it is essential to study the manufacturing behavior of the alloy. In recent years, additive manufacturing (AM) has gradually occupied an important position in material processing field, especially on machining complex three-dimensional near-net shaped components that cannot be manufactured by powder metallurgy, casting and deformation [6]. Selective laser melting (SLM) is a novel metal manufacturing method among the AM technologies. Because of its flexible layer-by-layer control strategy [7], SLM has received extensive attention on aerospace, biomedicine, energy motivation, etc. Compared with the traditional forging and casting process, SLM has a lot of advantages including energy saving, cost reduction, efficiency improvement and rapid preparation of complex shaped parts with high flexibility [8].
Corresponding author. State Key Laboratory of Solidification Processing Technology, Northwestern Polytechnical University, Xi'an, 710072, PR China. Corresponding author. E-mail addresses: [email protected] (R.D. Li), [email protected] (L. Huang).
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https://doi.org/10.1016/j.intermet.2018.10.018 Received 27 July 2018; Received in revised form 18 October 2018; Accepted 24 October 2018 0966-9795/ © 2018 Published by Elsevier Ltd.
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Meanwhile, due to its high melting and solidification rate (103–108 K/s) [9], unique microstructures such as ultrafine grain structure, metastable phases and supersaturated solid solution can be obtained by SLM [10]. SLM has been successfully applied to prepare Fe-based [11,12], Ti-based [13], Al-based [14,15], and amorphous alloy [16]; thus SLM holds great potential for fabrication of HEA with complex shape, ultrafine microstructure and high precision. At present, few people studied the fabrication of HEA by AM. Joseph et al. used laser deposition to fabricate Al0.3CrFeNi HEA and AlxCoCrFeNi, and discovered the tensile/compressive asymmetry [17,18]. Brief et al. studied the fabrication of FeCoCrNi HEA with SLM and concluded that this HEA presented a single solid solution, excellent tensile strength and elongation compared with other methods [19]. Similarly, some scholars reported the manufacturing of FeCoCrNiB, AlCoCrNiFe HEA by laser deposition or laser cladding [20,21]. Recently, R. Li et al. studied the fabrication of an equiatomic CoCrFeMnNi HEA by SLM, and found the little tetragonal σ precipitate phase, which is different from traditional methods [22]. In this paper, the fabrication of equimolar AlCoCrFeNi HEA by SLM was studied. As what is commonly known, not all the metal powder is suitable for SLM processing due to the fact that SLM is a rapid cooling and nonequilibrium metallurgical process, because that the thermal stress is inevitably generated in the printing process, resulting in the formation of solidification cracks, balls and pores. What's more, it is worth studying whether the AlCoCrFeNi HEA fabricated by SLM is still a fcc + bcc structure. Based on that, in the first time, HEA of AlCoCrFeNi is manufactured by SLM, with emphasis on microstructural evolution, phase identification and properties.
Table 1 The exact chemical composition of powder and the SLM samples. 68.4 J/mm3
83.3 J/mm3
97.2 J/mm3
111.1 J/mm3
Al Co Cr Fe Ni
10.5 21.1 19.9 22.4 26.1
10.3 21.1 19.3 22.2 27.1
10.1 21.5 19.4 22.0 27.0
9.98 21.4 19.6 22.2 22.2
10.1 21.6 19.7 22.0 26.6
3. Results and discussion 3.1. Densification of the AlCoCrFeNi Fig. 1 shows the effects of the VED on the density of the SLM processed AlCoCrFeNi samples. Obviously, the density of the SLM samples gradually increases as the VED rises to 97.2 J/mm3, beyond which the density doesn't further increase. The maximum relative density is 98.4%. The improvement of the relative density with the increased VED can be understood by the correlation between the viscosity and the VED. During SLM, the temperature field tends to increase when high VED is inputted into the powder bed. Relatively high VED (> 97.2 J/ mm3) speed results in the high temperature, which reduces weld pool tension and increases its flowability. The high temperature also allows for a sufficient amount of liquid phase diffusion and facilitates the interlayer bonding [24], and accordingly improves the densification behavior.
The powder material with the equimolar AlCoCrFeNi HEA was prepared by gas atomization, which was conducted in an inert and high purity argon atmosphere to avoid being oxidized. A laser particle size analyzer was applied to find the particle size of the powder (Ma-stersiga). The particle size of the powder ranges from 3.8 μm to 53 μm, and the average particle size is about 28.6 μm. The SLM samples were fabricated using a FS271 M machine (Farsoon, Inc, China), which was equipped with a 500 W fiber laser with the beam spot of 90 μm in diameter. During the SLM process, the samples were produced in an argon atmosphere with less than 0.1% of the oxygen content. A steel plate was used as the building substrate and the temperature of the steel plate was set as 100 °C. For starters, in order to obtain high density samples, different laser parameters were set. The scanning speed (v) is 1000 mm/s, and the laser power (P) ranges from 250 to 400 W. The layer thickness (t) is 0.04 mm, and the hatch spacing (h) is 0.09 mm. The scanning direction for the two consecutive layers N and N+1 is set as 67° during the SLM process. We have comprehensively considered the laser parameters (P, v, h, t) by volumetric energy density (VED) due to the layer-by-layer manner of SLM, which is defined as follows [23]:
P vht
Powder
(EBSD) were shown by scanning electron microscopy (SEM, FEI NanoLAB, 600i). The chemical compositions were analyzed by electron probe microanalyzer (EPMA, XAe8530, JEOL, Japan). Micro-hardness was tested by a micro-Vickers hardness tester with a load of 100 g for 15s (ASTM E 38-08). The nanoidentation test was carried out on a UNHT nanoindentation tester for the polished portion of the SLM samples and the loading-unloading test under 30 mN of load. The potentiodynamic polarization scanning was carried out at a scan rate of 1 mV per second.
2. Experimental procedures
VED =
Elements (wt. %)
3.2. Phase analysis Phase analysis of the powder and SLM processed samples are shown in Fig. 2. It can be found that all the printed samples consist of disordered body-center cubic phase (A2) and ordered body-center cubic phase (B2) as well as the original powder, Due to the fact that the gas atomization is also a rapid cooling process. It was verified that the B2
(1)
According to the formula above, the VED of samples by SLM is 69.4 J/mm3, 83.3 J/mm3, 97.2 J/mm3, and 111.1 J/mm3, respectively. The chemical composition of the original powder and SLM samples is measured based on inductively coupled plasma-atomic (ICP), as shown in Table 1. The Archimedes method was used to measure the relative density of each printed sample. Phase identification was performed in a X-ray Diffraction (XRD) with a D/max2500pc using a Cu-Ka radiation (K = 0.154 nm). Metallographic specimens were cut, ground and polished according to the standard procedure, and then etched with aqua regia for 3–5s. Microstructure observations were performed with a scanning electron microscope (SEM, Nova Nano SEM 230), and the energy dispersive spectrometer (EDS) was equipped. Images of crystal orientation and phase content by electron backscatter diffraction
Fig. 1. The effect of VED on the SLM processed density. 25
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Fig. 2. XRD pattern of AlCoCrFeNi powder and SLM processed samples: (a) different VED, (b) the peak of (110) enlarged drawing.
Fig. 3. Phase distribution of different VED of SLM samples: (a) 68.4 J/mm3; (b) 83.3 J/mm3; (c) 97.2 J/mm3; (a) 111.1 J/mm3.
increase of VED as shown in Fig. 3. The main reason was that the higher VED induced the higher cooling rate, which facilitated the generation of much more B2 phase.
and A2 phase in the AlCoCrFeNi HEA powder were retained during the SLM process. A. Munitz et al. studied the cast AlCoCrFeNi, and found that it contained the single bcc phase [25]. K.S.Lee studied the deformed AlCoCrFeNi, and found that it included the fcc + bcc dual phase [26]. These two results are different from that in this paper. However, it is observed that the peak of (110) becomes broader and is shifted to a lower 2θ angle (Fig. 3b) compared with the original powder with the increase of VED, suggesting that the SLM processing can produce ultra-fine grains due to the rapid cooling rate and the decreased inter-planar space. The decrease in lattice constant is because of the volatilization of Al element in the SLM process. It is known that Al element has the lowest boiling point (2327 °C) and the smallest vaporization heat (291.4 kJ/mol) compared with other elements. It also can be found that the relative intensities of (100) and (200) peak gradually increase with the VED, but the relative intensities of the (110) and (211) peak are in the opposite, indicating the much stronger (100) and (200) texture. It is well-known that the cooling rate has a great effect on the phase composition [27]. Previous studies have shown that the AlCoCrFeNi HEA contains single BCC or dual BCC + FCC or A2+B2 phases at different cooling rate [28–30]. Fig. 3 shows the phase distribution with different VEDs. It can be found that the B2 phase is mainly concentrated on the boundary of melt pool, so it can be speculated that the B2 phase is the primary phase. The main reason is that the combination between Al and Ni element is the strongest in the AlCoCrFeNi HEA [31]. Firstly, Al and Ni element are combined to form an ordered Al-Ni solid solution phase with BCC superlattice structure and the other elements are uniformly distributed around the Al-Ni phase (B2), forming a single A2 phase. What's more, metal liquid in the melt pool firstly solidifies on the boundary of the melt pool and has the fastest cooling rate during SLM. As the metallic liquid solidifies, the cooling rate is gradually reduced. More importantly, it can also be noticed that the A2 phase decreases while B2 phase increases with the
3.3. Microstructural evolution Fig. 4 shows the typical microstructure of SLM printed sample at the VED of 97.2 J/mm3. The boundary of the melt pool is clear (Fig. 4a). It can be found that bundles of grains grow along the same direction, indicating the epitaxial growth of the grains. In the process of the grain selection, the grain whose “easy-growth” direction is oriented along the direction of solidification out-grows. Therefore, the orientation of the solidification grain is usually perpendicular to the melt pool. SLM involves rapid solidification of melt pool and reciprocating heat cycle of multi melt pools, resulting in the elongated columnar grains which extend over multiple building layers. What's more, there exists obvious boundary of elongated columnar grain after the polished section is etched. Through characterizing the phase distribution by EBSD (Fig. 3), it can be confirmed that the phase on the boundary of the melt pool is the B2 phase. From Fig. 4b and c, it can be found that the formation of cellular structure is similar to the results proposed by Zhu et al. [32]. It is interesting that some precipitates are found in the sample without the heat treatment (Fig. 4d). To further identify the precipitate, the EDS maps are applied (Fig. 5). It can be found that the precipitate is rich in Fe and Cr (Fe-Cr phase). In the current work, the precipitate appears in the SLM-printed AlCoCrFeNi HEA, because that the SLM printed AlCoCrFeNi HEA has ultra-fine grains, which provides abundant grain boundaries, and then enhances the overall diffusion, facilitating the FeCr phase precipitation. To further analyze the distribution of the main alloy elements in macroscopic scale, the EPMA test showing the Al, Co, Cr, Fe, and Ni distribution was conducted, as what shown in Fig. 6. It can be found 26
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Fig. 4. SEM images showing the typical microstructure of SLM samples:(a) low mangnification, (b) is the high mangnification of (a); (c) and (d) are the high mangnification of (b). The dashed lines highlight the melt pool boundaries.
intensity increases and (110) texture intensity declines gradually in the SLM processed sample, which is consistent with the peak of XRD. The orientation distribution function (ODF) of SLM samples is calculated by the iterative series expansion method with the (111), (200), and (220) pole figures (Fig. 8). In the section of φ2 = 0°, there is no obvious orientation texture. In the section of φ2 = 45°, the texture is the strongest both in the side view and top view. The orientation density of top view and side view is 5.24 and 5.36 respectively, which means that the side view is more oriented than top view. With the increase of VED, the orientation density gradually increases from 5.24 to 5.34 in the top view of printed sample, and it can be found that the orthorhombic symmetry is weakened. What's more, it also can be found that the texture orientation is transformed from {112} < 1-12 > to {110} < 1-
that there is no evident segregation of the Al, Co, Cr, Fe, and Ni elements under SLM rapid solidification. Hence all the elements are uniformly distributed in the melt pool. 3.4. Texture analysis In order to observe the grain texture, the XRD macro-texture tests of SLM processed AlCoCrFeNi samples from the top views and side views were conducted. The pole figures (PF) of (200), (211) and (110) planes of SLM processed AlCoCrFeNi samples are shown in Fig. 7. When the VED is 68.4 J/mm3, weak (200) texture is found in both side view and top view, and there is a more intensified (200) orientation in the top view than the side view. As the VED rises to 111.1 J/mm3, (200) texture
Fig. 5. EDS maps showing the elements distribution in cellular grain (sub-micro scale). 27
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Fig. 6. EPMA maps of the SLM samples showing elements distribution at multi melt pools (200 μm scale).
process of constant remelting and solidification. Normally, the columnar grain originates from the first several building layers. The first several layers have been identified to contain randomly oriented equiaxed grains due to the rapid heat loss in the building substrate and the radiation from the melt pool surface [33]. As the building process goes on, the columnar grains begin to grow epitaxially in the material's
12 > with the increase of VED. In order to further observe the grain texture, the EBSD images of SLM processed AlCoCrFeNi samples in side and top views are provided in Fig. 9. In the side view, the grains grow epitaxially and are transformed into columnar grains (Fig. 9a–d), and the columnar grains extend over multiple building layers. The main reason is that SLM is a
Fig. 7. XRD Pole figure (PF) of SLM printed samples with different VED: (a) and (b) 68.4 J/mm3, (c) 97.2 J/mm3, (d) 111.1 J/mm3. 28
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Fig. 8. Orientation distribution function (ODF) of SLM printed samples with different VED: (a) and (b) are the standard of bcc structure in 0°and 45°; (c) the top view of68.4 J/mm3; (d) the side view of 68.4 J/mm3, (e) the side view of 97.2 J/mm3, (f) the side view of 111.1 J/mm3.
order to further analyze the orientation relationship between A2 and B2 phase from the PF (Fig. 10), it can be found that the PF of {100} of A2 phase and {100} of B2 phase in the red dotted line have the same point, indicating that there is a {100}A2∥{100}B2 orientation relationship. The formation of A2 and B2 phase has a great relationship with the cooling rate. At a high cooling rate, the B2 phase is firstly formed. As the solidification process goes on, the cooling rate gradually decreases and then the A2 phase is formed. Due to the difference in the cooling rate, different phases are formed one after another and a specific orientation relationship is generally formed.
preferred orientation. In the top view (Fig. 9e), It can be clearly seen that the scanning track is composed of columnar grains inside the scanning track, while the boundary of the scanning track is composed of equiaxed grains. Due to the fact that laser energy has a Gaussian distribution and the laser beam moves rapidly during the scanning process, the temperature gradient inside the scan track is large, and there are different cooling rates in different parts of the scan track, and the cooling rate reaches its maximum at the center, and then the cooling rate gradually decreases from the center to the boundary. Therefore, the grain size in the center of the scan track is minimized, and the grain size on its boundary is maximized. Meanwhile, the scanning track with a width of 90 μm (same as the hatching distance) shows the grating pattern of the laser beam on the powder surface. In the side view, from the pole figures in the EBSD images (Fig. 9a–d), it can be figured out that there is a great change in grain orientation (grain colour), and the main grain colors are red, blue and green, indicating that the grain orientation is (001), (101) and (111). The main reason is that columnar grains are often hindered at the solidification front, inducing the formation of new grains with random crystallographic orientation and more evident isotropic characters. The microstructure was dominated by thermal gradients and solidification rates, which are related to the parameters. In order to further observe the grain texture and orientation relationship between different phases with different VED, the Inverse Pole Figures (IPF) and Pole Figures (PF) are shown in Fig. 10. It can be found that the SLM samples have obvious preferred orientation in < 001 > crystallographic direction families at a low VED, and is gradually transformed into < 111 > crystallographic direction families at a higher VED. This is because that the VED has a significant influence on the shape of melt pool, and the epitaxial growth of the grains is perpendicular to the boundary of melt pool. Therefore, the VED has an evident effect on the preferred orientation of grains. In
3.5. Micro-hardness of different VED of SLM samples Fig. 11 shows the micro-hardness of SLM samples fabricated with various VEDs. It is obvious that the micro-hardness increases significantly with the VED. When the VED is 111.1 J/mm3, the maximum micro-hardness is 632.8Hv, which is much higher than that of the sample fabricated by other methods, such as laser engineered net shaping (LENS) [34], elective electron beam melting (SEBM) [35] and casting [36]. Significantly refined grains can be formed and elements are distributed uniformly due to the fact that SLM is a rapid cooling process [37,38]. It also can be found that the micro-hardness increases with the VED. The main reason is that higher VED leads to the larger cooling rate and smaller grain size, which is consistent with the HallPetch theory. Meanwhile, a large number of dislocations can be created by the excessive cooling rate. According to the dislocation theory, the grain boundary hinders the dislocation movement, and the grain boundary expands when the grains are refined. Therefore, larger stress is required for the dislocation movement and grain deformation, which indicates a higher resistance to external loads (in other words, microhardness increases). In addition, high-density dislocations induce 29
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Fig. 9. EBSD images pole figure (IPF) of SLM printed samples with different VED: (a) 68.4 J/mm3, (b) 83.3 J/mm3, (c) 97.2 J/mm3, (d) and (e) 111.1 J/mm3.
phase (a hard and brittle phase). This is another reason why the microhardness of the SLM AlCoCrFeNi samples is higher than that of the samples fabricated traditional powder metallurgy. Fig. 12 shows the load-displacement curves of SLM samples with
dislocation packing and entanglements, and hinder grain boundary sliding and dislocation movement, which also helps improve the microhardness. What's more, the existence of disordered B2 phase will also increase its micro-hardness due to the fact that B2 phase is a Ni-Al
Fig. 10. Shows Inverse Pole Figure (IPF) and Pole Figures (PF) of different VED and phase of SLM samples. 30
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Table 2 The indentation depths, microhardness and Young's modulus of SLM samples with different VED. VED (J/mm3) Hmax (nm) Nanohardness (HV) E (GPa)
111.1 404.66 853.23 198.90
97.2 418.61 776.25 189.18
83.3 427.19 740.66 185.24
68.4 451.39 722.40 184.95
Fig. 11. Microhardness of SLM samples fabricated with different VEDs and methods.
different VEDs through nano-indentation test. The indentation depths of the SLM samples gradually increase with the decrease of VED (Fig. 12a). On the contrary, it can be observed that the nanohardness and Young's modulus of the sample decrease with the decrease of VED, as revealed in Table 1. Owing to the higher cooling rate with the increase of VED during the SLM process, finer microstructure is found in the samples, which can further increase the nanohardness and Young's modulus. As we all know, the Young's modulus is sensitive to the phase composition. The dual AlCoCrFeNi HEA consists of A2 and B2 phase, and the B2 is hard brittle phase. So the increase of B2 leads to the higher Young's modulus. What's more, residual stress is very large during the additive manufacturing process, which is advantageous to the improvement of nanohardness and Young's modulus [39]. Appropriate residual stress of the SLM samples helps increase the hardness [40] when the SLM samples reach a sufficiently high densification and no micro-cracks or micro-holes are found (see Table 2). The Tafel curves for SLM samples with different VEDs in aerated 3.5% NaCl solutions are shown in Fig. 13. Corrosion potentials and currents can be detected from the Tafel curves, reflecting the corrosion resistance. It can be found that the VED has a significant effect on the corrosion resistance. The corrosion resistance is improved with the increase of VED, reflecting the potential increase. Owing to the fast cooling rate, a finer microstructure and more B2 phases are formed. Because that the B2 (Ni-Al) phase are mainly distributed along the columnar grains, it has the best corrosion resistance. Thus the increase of B2 phase content helps increase the corrosion resistance of the SLM samples. What's more, the lack of anode and cathode active surfaces
Fig. 13. Potentiodynamic polarization curves of the SLM samples with different VED.
reduces the susceptibility to corrosion in the melt pool, which further improves the corrosion resistance. 4. Conclusions In this paper, the equimolar AlCoCrFeNi HEA powder was printed by SLM, with emphasize on its densification, phase identify, microstructure evolution, crystallographic texture, micro-hardness and correction resistance. The following conclusions are drawn: 1) The density of the SLM samples gradually increases as the VED rises up to 97.2 J/mm3, beyond which the density doesn't increase further and the maximum relative density is 98.4%. 2) The SLM printed AlCoCrFeNi HEA consists of dual A2 and B2 phases, which is different from the conventional cast and deformation processed samples with A2 and fcc phases due to the rapid cooling rate of SLM. Besides, the content of B2 phase increases with the enhanced VED in the SLM sample. 3) The microstructure of SLM samples shows the epitaxial growth of A2
Fig. 12. (a) Load-displacement curves of the nanoindentation test performed of SLM samples with different VED, (b)–(e) microphotograph of nanoidentation. 31
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grain bundles perpendicular to the melt pool boundary; the B2 phase is between the columnar A2 grains. Interestingly, the Fe-Cr precipitate phase is found in the sample which did not experience the heat treatment. 4) The texture of the SLM sample has apparent relation with VED, the SLM sample shows < 001 > preferred orientation at low VED, and the orientation is gradually transformed into < 111 > crystallographic direction families at a higher VED. Besides, the orientation relationship between A2 and B2 phases is {100}A2∥{100}B2. 5) The micro-hardness and the corrosion resistance are gradually improved as the VED rises and the maximum micro-hardness rises to 632.8HV, which is much higher than that of the samples fabricated by other methods (> 100Hv). The main reason is that the higher VED results in the much more rapid cooling rate and the formation of more B2 phases.
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Acknowledgements The authors wish to acknowledge the financial support of National Key R&D Program of China (No. 2017YFB0306305, 2016YFB1100101), National Natural Science Foundation of China (51871249, 51505166), Research fund of the State Key Laboratory of Solidification Processing of NWPU (SKLSP201748), Science and Technology Planning Project of Changsha, Science and Technology Planning Project of Shenzhen, Guizhou Science and Technology Plan Project (grant No. GHK support [2017]2013) and Huxiang Young talents (2018RS3007). References [1] B. Cantor, I. Chang, P. Knight, A. Vincent, Microstructural development in equiatomic multicomponent alloys, Mat. Sci. Eng. a Struct. 375 (2004) 213–218. [2] Y.F. Ye, Q. Wang, J. Lu, C.T. Liu, Y. Yang, High-entropy alloy: challenges and prospects, Mater. Today 19 (2016) 349–362. [3] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (2004) 299–303. [4] Y. Zhang, Mechanical properties and structures of high entropy alloys and bulk metallic glasses composites, in: J.F. Nie, A. Morton (Eds.), Materials Science Forum, 2010, pp. 1058–1061. [5] M.J. Jang, S. Praveen, H.J. Sung, J.W. Bae, J. Moon, H.S. Kim, High-temperature tensile deformation behavior of hot rolled CrMnFeCoNi high-entropy alloy, J. Alloy. Comp. 730 (2018) 242–248. [6] T. DebRoy, H.L. Wei, J.S. Zuback, T. Mukherjee, J.W. Elmer, J.O. Milewski, A.M. Beese, A. Wilson-Heid, A. De, W. Zhang, Additive manufacturing of metallic components - process, structure and properties, Prog. Mater. Sci. 92 (2018) 112–224. [7] 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) 401–477. [8] H. Rao, S. Giet, K. Yang, X. Wu, C.H.J. Davies, The influence of processing parameters on aluminium alloy A357 manufactured by Selective Laser Melting, Mater. Des. 109 (2016) 334–346. [9] Y. Li, K. Zhou, P. Tan, S.B. Tor, C.K. Chua, K.F. Leong, Modeling temperature and residual stress fields in selective laser melting, Int. J. Mech. Sci. 136 (2018) 24–35. [10] X.P. Li, X.J. Wang, M. Saunders, A. Suvorova, L.C. Zhang, Y.J. Liu, M.H. Fang, Z.H. Huang, T.B. Sercombe, A selective laser melting and solution heat treatment refined Al-12Si alloy with a controllable ultrafine eutectic microstructure and 25% tensile ductility, Acta Mater. 95 (2015) 74–82. [11] R. Li, Y. Shi, Z. Wang, L. Wang, J. Liu, W. Jiang, Densification behavior of gas and water atomized 316L stainless steel powder during selective laser melting, Appl. Surf. Sci. 256 (2010) 4350–4356. [12] B. Song, Z. Wang, Q. Yan, Y. Zhang, J. Zhang, C. Cai, Q. Wei, Y. Shi, Integral method of preparation and fabrication of metal matrix composite: selective laser melting of in-situ nano/submicro-sized carbides reinforced iron matrix composites, Mat. Sci. Eng. a Struct. 707 (2017) 478–487. [13] L. Zhou, T. Yuan, R. Li, J. Tang, M. Wang, F. Mei, Anisotropic mechanical behavior of biomedical Ti-13Nb-13Zr alloy manufactured by selective laser melting, J. Alloy. Comp. 762 (2018) 289–300. [14] R. Li, M. Wang, T. Yuan, B. Song, C. Chen, K. 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) 117–128.
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Microstructures and properties of an equimolar AlCoCrFeNi high entropy alloy printed by selective laser melting
T
P.D. Niua, R.D. Lia,b,∗, T.C. Yuana, S.Y. Zhuc, C. Chena, M.B. Wanga, L. Huanga,∗∗ a
State Key Laboratory of Powder Metallurgy, Science and Technology on High Strength Structural Materials Laboratory, Central South University, Changsha, 410083, PR China b State Key Laboratory of Solidification Processing Technology, Northwestern Polytechnical University, Xi'an, 710072, PR China c Materials Institute of Genome, Shanghai University, Shanghai, 200444, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: High entropy alloys Selective laser melting Equiatomic AlCoCrFeNi
The equimolar AlCoCrFeNi high entropy alloy (HEA) was printed by selective laser melting (SLM), with emphasis on its densification, phase identification, non-equilibrium microstructure and properties. The density of SLM samples increase gradually with the increase of volumetric energy density (VED) with the maximal relative density 98.4%. SLM samples are consisted of disordered body-center cubic phases (A2) and ordered body-center cubic phases (B2) which are different from the conventional cast or deformation with A2 and fcc phase; the content of B2 phase increases with the VED. Interestingly, the Fe-Cr precipitate, which has never been found in the cast one, is observed in the SLM one without heat treatment. The microstructure of SLM samples shows epitaxial growth of columnar A2 grain bundles perpendicular to the melt pool boundary with an average grain size of about 1.5 μm; the B2 phase is between the columnar A2 grains. At low VED, the SLM sample shows < 001 > preferred orientation, and is gradually transformed to the < 111 > crystallographic direction families at a higher VED. Besides, there exists the {100}A2∥{100}B2 orientation relationship in A2 and B2 phase. The maximal micro-hardness of 632.8Hv of SLM sample is higher than the cast one and the electron beam melting printed one.
1. Introduction The development of new materials is very important for the material science and the ecological utilization of limited resources on the earth. Conventional alloys like Al- alloy, Fe-alloy, Ti-alloys, TiAl-alloys, FeAlalloys are based on one or two main elements, and a relatively small amount of other elements are added to modify the microstructure and properties of the pure metal base. However, as a new class of metal alloy, high entropy alloys (HEAs) which provide new strategies and expand the range of traditional alloys for alloy design [1] attract widespread attention in recent years due to its excellent mechanical, high thermal stability, and outstanding wear resistance [2]. The concept of HEA was first proposed by Yeh et al. in 2004 [3]. In most of the cases, HEAs which are defined as alloys containing five or more major alloying elements possess a concentration between 5 at% and 35 at%, forming a single or double phase crystalline structure despite that there are multiple elements with different crystal structures. In particular, the quinary HEA composed of Al, Co, Cr, Fe, and Ni elements has received
∗
much attention due to its excellent properties such as good corrosion resistance, high hardness and high yield strength even at high temperatures [4]. Usually, HEA is manufactured by conventional casting or plastic deformation [5]; however, these methods fail to machine complex shapes and ultrafine grained structure. Therefore, for a HEA which has a promising application, it is essential to study the manufacturing behavior of the alloy. In recent years, additive manufacturing (AM) has gradually occupied an important position in material processing field, especially on machining complex three-dimensional near-net shaped components that cannot be manufactured by powder metallurgy, casting and deformation [6]. Selective laser melting (SLM) is a novel metal manufacturing method among the AM technologies. Because of its flexible layer-by-layer control strategy [7], SLM has received extensive attention on aerospace, biomedicine, energy motivation, etc. Compared with the traditional forging and casting process, SLM has a lot of advantages including energy saving, cost reduction, efficiency improvement and rapid preparation of complex shaped parts with high flexibility [8].
Corresponding author. State Key Laboratory of Solidification Processing Technology, Northwestern Polytechnical University, Xi'an, 710072, PR China. Corresponding author. E-mail addresses: [email protected] (R.D. Li), [email protected] (L. Huang).
∗∗
https://doi.org/10.1016/j.intermet.2018.10.018 Received 27 July 2018; Received in revised form 18 October 2018; Accepted 24 October 2018 0966-9795/ © 2018 Published by Elsevier Ltd.
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Meanwhile, due to its high melting and solidification rate (103–108 K/s) [9], unique microstructures such as ultrafine grain structure, metastable phases and supersaturated solid solution can be obtained by SLM [10]. SLM has been successfully applied to prepare Fe-based [11,12], Ti-based [13], Al-based [14,15], and amorphous alloy [16]; thus SLM holds great potential for fabrication of HEA with complex shape, ultrafine microstructure and high precision. At present, few people studied the fabrication of HEA by AM. Joseph et al. used laser deposition to fabricate Al0.3CrFeNi HEA and AlxCoCrFeNi, and discovered the tensile/compressive asymmetry [17,18]. Brief et al. studied the fabrication of FeCoCrNi HEA with SLM and concluded that this HEA presented a single solid solution, excellent tensile strength and elongation compared with other methods [19]. Similarly, some scholars reported the manufacturing of FeCoCrNiB, AlCoCrNiFe HEA by laser deposition or laser cladding [20,21]. Recently, R. Li et al. studied the fabrication of an equiatomic CoCrFeMnNi HEA by SLM, and found the little tetragonal σ precipitate phase, which is different from traditional methods [22]. In this paper, the fabrication of equimolar AlCoCrFeNi HEA by SLM was studied. As what is commonly known, not all the metal powder is suitable for SLM processing due to the fact that SLM is a rapid cooling and nonequilibrium metallurgical process, because that the thermal stress is inevitably generated in the printing process, resulting in the formation of solidification cracks, balls and pores. What's more, it is worth studying whether the AlCoCrFeNi HEA fabricated by SLM is still a fcc + bcc structure. Based on that, in the first time, HEA of AlCoCrFeNi is manufactured by SLM, with emphasis on microstructural evolution, phase identification and properties.
Table 1 The exact chemical composition of powder and the SLM samples. 68.4 J/mm3
83.3 J/mm3
97.2 J/mm3
111.1 J/mm3
Al Co Cr Fe Ni
10.5 21.1 19.9 22.4 26.1
10.3 21.1 19.3 22.2 27.1
10.1 21.5 19.4 22.0 27.0
9.98 21.4 19.6 22.2 22.2
10.1 21.6 19.7 22.0 26.6
3. Results and discussion 3.1. Densification of the AlCoCrFeNi Fig. 1 shows the effects of the VED on the density of the SLM processed AlCoCrFeNi samples. Obviously, the density of the SLM samples gradually increases as the VED rises to 97.2 J/mm3, beyond which the density doesn't further increase. The maximum relative density is 98.4%. The improvement of the relative density with the increased VED can be understood by the correlation between the viscosity and the VED. During SLM, the temperature field tends to increase when high VED is inputted into the powder bed. Relatively high VED (> 97.2 J/ mm3) speed results in the high temperature, which reduces weld pool tension and increases its flowability. The high temperature also allows for a sufficient amount of liquid phase diffusion and facilitates the interlayer bonding [24], and accordingly improves the densification behavior.
The powder material with the equimolar AlCoCrFeNi HEA was prepared by gas atomization, which was conducted in an inert and high purity argon atmosphere to avoid being oxidized. A laser particle size analyzer was applied to find the particle size of the powder (Ma-stersiga). The particle size of the powder ranges from 3.8 μm to 53 μm, and the average particle size is about 28.6 μm. The SLM samples were fabricated using a FS271 M machine (Farsoon, Inc, China), which was equipped with a 500 W fiber laser with the beam spot of 90 μm in diameter. During the SLM process, the samples were produced in an argon atmosphere with less than 0.1% of the oxygen content. A steel plate was used as the building substrate and the temperature of the steel plate was set as 100 °C. For starters, in order to obtain high density samples, different laser parameters were set. The scanning speed (v) is 1000 mm/s, and the laser power (P) ranges from 250 to 400 W. The layer thickness (t) is 0.04 mm, and the hatch spacing (h) is 0.09 mm. The scanning direction for the two consecutive layers N and N+1 is set as 67° during the SLM process. We have comprehensively considered the laser parameters (P, v, h, t) by volumetric energy density (VED) due to the layer-by-layer manner of SLM, which is defined as follows [23]:
P vht
Powder
(EBSD) were shown by scanning electron microscopy (SEM, FEI NanoLAB, 600i). The chemical compositions were analyzed by electron probe microanalyzer (EPMA, XAe8530, JEOL, Japan). Micro-hardness was tested by a micro-Vickers hardness tester with a load of 100 g for 15s (ASTM E 38-08). The nanoidentation test was carried out on a UNHT nanoindentation tester for the polished portion of the SLM samples and the loading-unloading test under 30 mN of load. The potentiodynamic polarization scanning was carried out at a scan rate of 1 mV per second.
2. Experimental procedures
VED =
Elements (wt. %)
3.2. Phase analysis Phase analysis of the powder and SLM processed samples are shown in Fig. 2. It can be found that all the printed samples consist of disordered body-center cubic phase (A2) and ordered body-center cubic phase (B2) as well as the original powder, Due to the fact that the gas atomization is also a rapid cooling process. It was verified that the B2
(1)
According to the formula above, the VED of samples by SLM is 69.4 J/mm3, 83.3 J/mm3, 97.2 J/mm3, and 111.1 J/mm3, respectively. The chemical composition of the original powder and SLM samples is measured based on inductively coupled plasma-atomic (ICP), as shown in Table 1. The Archimedes method was used to measure the relative density of each printed sample. Phase identification was performed in a X-ray Diffraction (XRD) with a D/max2500pc using a Cu-Ka radiation (K = 0.154 nm). Metallographic specimens were cut, ground and polished according to the standard procedure, and then etched with aqua regia for 3–5s. Microstructure observations were performed with a scanning electron microscope (SEM, Nova Nano SEM 230), and the energy dispersive spectrometer (EDS) was equipped. Images of crystal orientation and phase content by electron backscatter diffraction
Fig. 1. The effect of VED on the SLM processed density. 25
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Fig. 2. XRD pattern of AlCoCrFeNi powder and SLM processed samples: (a) different VED, (b) the peak of (110) enlarged drawing.
Fig. 3. Phase distribution of different VED of SLM samples: (a) 68.4 J/mm3; (b) 83.3 J/mm3; (c) 97.2 J/mm3; (a) 111.1 J/mm3.
increase of VED as shown in Fig. 3. The main reason was that the higher VED induced the higher cooling rate, which facilitated the generation of much more B2 phase.
and A2 phase in the AlCoCrFeNi HEA powder were retained during the SLM process. A. Munitz et al. studied the cast AlCoCrFeNi, and found that it contained the single bcc phase [25]. K.S.Lee studied the deformed AlCoCrFeNi, and found that it included the fcc + bcc dual phase [26]. These two results are different from that in this paper. However, it is observed that the peak of (110) becomes broader and is shifted to a lower 2θ angle (Fig. 3b) compared with the original powder with the increase of VED, suggesting that the SLM processing can produce ultra-fine grains due to the rapid cooling rate and the decreased inter-planar space. The decrease in lattice constant is because of the volatilization of Al element in the SLM process. It is known that Al element has the lowest boiling point (2327 °C) and the smallest vaporization heat (291.4 kJ/mol) compared with other elements. It also can be found that the relative intensities of (100) and (200) peak gradually increase with the VED, but the relative intensities of the (110) and (211) peak are in the opposite, indicating the much stronger (100) and (200) texture. It is well-known that the cooling rate has a great effect on the phase composition [27]. Previous studies have shown that the AlCoCrFeNi HEA contains single BCC or dual BCC + FCC or A2+B2 phases at different cooling rate [28–30]. Fig. 3 shows the phase distribution with different VEDs. It can be found that the B2 phase is mainly concentrated on the boundary of melt pool, so it can be speculated that the B2 phase is the primary phase. The main reason is that the combination between Al and Ni element is the strongest in the AlCoCrFeNi HEA [31]. Firstly, Al and Ni element are combined to form an ordered Al-Ni solid solution phase with BCC superlattice structure and the other elements are uniformly distributed around the Al-Ni phase (B2), forming a single A2 phase. What's more, metal liquid in the melt pool firstly solidifies on the boundary of the melt pool and has the fastest cooling rate during SLM. As the metallic liquid solidifies, the cooling rate is gradually reduced. More importantly, it can also be noticed that the A2 phase decreases while B2 phase increases with the
3.3. Microstructural evolution Fig. 4 shows the typical microstructure of SLM printed sample at the VED of 97.2 J/mm3. The boundary of the melt pool is clear (Fig. 4a). It can be found that bundles of grains grow along the same direction, indicating the epitaxial growth of the grains. In the process of the grain selection, the grain whose “easy-growth” direction is oriented along the direction of solidification out-grows. Therefore, the orientation of the solidification grain is usually perpendicular to the melt pool. SLM involves rapid solidification of melt pool and reciprocating heat cycle of multi melt pools, resulting in the elongated columnar grains which extend over multiple building layers. What's more, there exists obvious boundary of elongated columnar grain after the polished section is etched. Through characterizing the phase distribution by EBSD (Fig. 3), it can be confirmed that the phase on the boundary of the melt pool is the B2 phase. From Fig. 4b and c, it can be found that the formation of cellular structure is similar to the results proposed by Zhu et al. [32]. It is interesting that some precipitates are found in the sample without the heat treatment (Fig. 4d). To further identify the precipitate, the EDS maps are applied (Fig. 5). It can be found that the precipitate is rich in Fe and Cr (Fe-Cr phase). In the current work, the precipitate appears in the SLM-printed AlCoCrFeNi HEA, because that the SLM printed AlCoCrFeNi HEA has ultra-fine grains, which provides abundant grain boundaries, and then enhances the overall diffusion, facilitating the FeCr phase precipitation. To further analyze the distribution of the main alloy elements in macroscopic scale, the EPMA test showing the Al, Co, Cr, Fe, and Ni distribution was conducted, as what shown in Fig. 6. It can be found 26
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Fig. 4. SEM images showing the typical microstructure of SLM samples:(a) low mangnification, (b) is the high mangnification of (a); (c) and (d) are the high mangnification of (b). The dashed lines highlight the melt pool boundaries.
intensity increases and (110) texture intensity declines gradually in the SLM processed sample, which is consistent with the peak of XRD. The orientation distribution function (ODF) of SLM samples is calculated by the iterative series expansion method with the (111), (200), and (220) pole figures (Fig. 8). In the section of φ2 = 0°, there is no obvious orientation texture. In the section of φ2 = 45°, the texture is the strongest both in the side view and top view. The orientation density of top view and side view is 5.24 and 5.36 respectively, which means that the side view is more oriented than top view. With the increase of VED, the orientation density gradually increases from 5.24 to 5.34 in the top view of printed sample, and it can be found that the orthorhombic symmetry is weakened. What's more, it also can be found that the texture orientation is transformed from {112} < 1-12 > to {110} < 1-
that there is no evident segregation of the Al, Co, Cr, Fe, and Ni elements under SLM rapid solidification. Hence all the elements are uniformly distributed in the melt pool. 3.4. Texture analysis In order to observe the grain texture, the XRD macro-texture tests of SLM processed AlCoCrFeNi samples from the top views and side views were conducted. The pole figures (PF) of (200), (211) and (110) planes of SLM processed AlCoCrFeNi samples are shown in Fig. 7. When the VED is 68.4 J/mm3, weak (200) texture is found in both side view and top view, and there is a more intensified (200) orientation in the top view than the side view. As the VED rises to 111.1 J/mm3, (200) texture
Fig. 5. EDS maps showing the elements distribution in cellular grain (sub-micro scale). 27
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Fig. 6. EPMA maps of the SLM samples showing elements distribution at multi melt pools (200 μm scale).
process of constant remelting and solidification. Normally, the columnar grain originates from the first several building layers. The first several layers have been identified to contain randomly oriented equiaxed grains due to the rapid heat loss in the building substrate and the radiation from the melt pool surface [33]. As the building process goes on, the columnar grains begin to grow epitaxially in the material's
12 > with the increase of VED. In order to further observe the grain texture, the EBSD images of SLM processed AlCoCrFeNi samples in side and top views are provided in Fig. 9. In the side view, the grains grow epitaxially and are transformed into columnar grains (Fig. 9a–d), and the columnar grains extend over multiple building layers. The main reason is that SLM is a
Fig. 7. XRD Pole figure (PF) of SLM printed samples with different VED: (a) and (b) 68.4 J/mm3, (c) 97.2 J/mm3, (d) 111.1 J/mm3. 28
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Fig. 8. Orientation distribution function (ODF) of SLM printed samples with different VED: (a) and (b) are the standard of bcc structure in 0°and 45°; (c) the top view of68.4 J/mm3; (d) the side view of 68.4 J/mm3, (e) the side view of 97.2 J/mm3, (f) the side view of 111.1 J/mm3.
order to further analyze the orientation relationship between A2 and B2 phase from the PF (Fig. 10), it can be found that the PF of {100} of A2 phase and {100} of B2 phase in the red dotted line have the same point, indicating that there is a {100}A2∥{100}B2 orientation relationship. The formation of A2 and B2 phase has a great relationship with the cooling rate. At a high cooling rate, the B2 phase is firstly formed. As the solidification process goes on, the cooling rate gradually decreases and then the A2 phase is formed. Due to the difference in the cooling rate, different phases are formed one after another and a specific orientation relationship is generally formed.
preferred orientation. In the top view (Fig. 9e), It can be clearly seen that the scanning track is composed of columnar grains inside the scanning track, while the boundary of the scanning track is composed of equiaxed grains. Due to the fact that laser energy has a Gaussian distribution and the laser beam moves rapidly during the scanning process, the temperature gradient inside the scan track is large, and there are different cooling rates in different parts of the scan track, and the cooling rate reaches its maximum at the center, and then the cooling rate gradually decreases from the center to the boundary. Therefore, the grain size in the center of the scan track is minimized, and the grain size on its boundary is maximized. Meanwhile, the scanning track with a width of 90 μm (same as the hatching distance) shows the grating pattern of the laser beam on the powder surface. In the side view, from the pole figures in the EBSD images (Fig. 9a–d), it can be figured out that there is a great change in grain orientation (grain colour), and the main grain colors are red, blue and green, indicating that the grain orientation is (001), (101) and (111). The main reason is that columnar grains are often hindered at the solidification front, inducing the formation of new grains with random crystallographic orientation and more evident isotropic characters. The microstructure was dominated by thermal gradients and solidification rates, which are related to the parameters. In order to further observe the grain texture and orientation relationship between different phases with different VED, the Inverse Pole Figures (IPF) and Pole Figures (PF) are shown in Fig. 10. It can be found that the SLM samples have obvious preferred orientation in < 001 > crystallographic direction families at a low VED, and is gradually transformed into < 111 > crystallographic direction families at a higher VED. This is because that the VED has a significant influence on the shape of melt pool, and the epitaxial growth of the grains is perpendicular to the boundary of melt pool. Therefore, the VED has an evident effect on the preferred orientation of grains. In
3.5. Micro-hardness of different VED of SLM samples Fig. 11 shows the micro-hardness of SLM samples fabricated with various VEDs. It is obvious that the micro-hardness increases significantly with the VED. When the VED is 111.1 J/mm3, the maximum micro-hardness is 632.8Hv, which is much higher than that of the sample fabricated by other methods, such as laser engineered net shaping (LENS) [34], elective electron beam melting (SEBM) [35] and casting [36]. Significantly refined grains can be formed and elements are distributed uniformly due to the fact that SLM is a rapid cooling process [37,38]. It also can be found that the micro-hardness increases with the VED. The main reason is that higher VED leads to the larger cooling rate and smaller grain size, which is consistent with the HallPetch theory. Meanwhile, a large number of dislocations can be created by the excessive cooling rate. According to the dislocation theory, the grain boundary hinders the dislocation movement, and the grain boundary expands when the grains are refined. Therefore, larger stress is required for the dislocation movement and grain deformation, which indicates a higher resistance to external loads (in other words, microhardness increases). In addition, high-density dislocations induce 29
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Fig. 9. EBSD images pole figure (IPF) of SLM printed samples with different VED: (a) 68.4 J/mm3, (b) 83.3 J/mm3, (c) 97.2 J/mm3, (d) and (e) 111.1 J/mm3.
phase (a hard and brittle phase). This is another reason why the microhardness of the SLM AlCoCrFeNi samples is higher than that of the samples fabricated traditional powder metallurgy. Fig. 12 shows the load-displacement curves of SLM samples with
dislocation packing and entanglements, and hinder grain boundary sliding and dislocation movement, which also helps improve the microhardness. What's more, the existence of disordered B2 phase will also increase its micro-hardness due to the fact that B2 phase is a Ni-Al
Fig. 10. Shows Inverse Pole Figure (IPF) and Pole Figures (PF) of different VED and phase of SLM samples. 30
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Table 2 The indentation depths, microhardness and Young's modulus of SLM samples with different VED. VED (J/mm3) Hmax (nm) Nanohardness (HV) E (GPa)
111.1 404.66 853.23 198.90
97.2 418.61 776.25 189.18
83.3 427.19 740.66 185.24
68.4 451.39 722.40 184.95
Fig. 11. Microhardness of SLM samples fabricated with different VEDs and methods.
different VEDs through nano-indentation test. The indentation depths of the SLM samples gradually increase with the decrease of VED (Fig. 12a). On the contrary, it can be observed that the nanohardness and Young's modulus of the sample decrease with the decrease of VED, as revealed in Table 1. Owing to the higher cooling rate with the increase of VED during the SLM process, finer microstructure is found in the samples, which can further increase the nanohardness and Young's modulus. As we all know, the Young's modulus is sensitive to the phase composition. The dual AlCoCrFeNi HEA consists of A2 and B2 phase, and the B2 is hard brittle phase. So the increase of B2 leads to the higher Young's modulus. What's more, residual stress is very large during the additive manufacturing process, which is advantageous to the improvement of nanohardness and Young's modulus [39]. Appropriate residual stress of the SLM samples helps increase the hardness [40] when the SLM samples reach a sufficiently high densification and no micro-cracks or micro-holes are found (see Table 2). The Tafel curves for SLM samples with different VEDs in aerated 3.5% NaCl solutions are shown in Fig. 13. Corrosion potentials and currents can be detected from the Tafel curves, reflecting the corrosion resistance. It can be found that the VED has a significant effect on the corrosion resistance. The corrosion resistance is improved with the increase of VED, reflecting the potential increase. Owing to the fast cooling rate, a finer microstructure and more B2 phases are formed. Because that the B2 (Ni-Al) phase are mainly distributed along the columnar grains, it has the best corrosion resistance. Thus the increase of B2 phase content helps increase the corrosion resistance of the SLM samples. What's more, the lack of anode and cathode active surfaces
Fig. 13. Potentiodynamic polarization curves of the SLM samples with different VED.
reduces the susceptibility to corrosion in the melt pool, which further improves the corrosion resistance. 4. Conclusions In this paper, the equimolar AlCoCrFeNi HEA powder was printed by SLM, with emphasize on its densification, phase identify, microstructure evolution, crystallographic texture, micro-hardness and correction resistance. The following conclusions are drawn: 1) The density of the SLM samples gradually increases as the VED rises up to 97.2 J/mm3, beyond which the density doesn't increase further and the maximum relative density is 98.4%. 2) The SLM printed AlCoCrFeNi HEA consists of dual A2 and B2 phases, which is different from the conventional cast and deformation processed samples with A2 and fcc phases due to the rapid cooling rate of SLM. Besides, the content of B2 phase increases with the enhanced VED in the SLM sample. 3) The microstructure of SLM samples shows the epitaxial growth of A2
Fig. 12. (a) Load-displacement curves of the nanoindentation test performed of SLM samples with different VED, (b)–(e) microphotograph of nanoidentation. 31
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grain bundles perpendicular to the melt pool boundary; the B2 phase is between the columnar A2 grains. Interestingly, the Fe-Cr precipitate phase is found in the sample which did not experience the heat treatment. 4) The texture of the SLM sample has apparent relation with VED, the SLM sample shows < 001 > preferred orientation at low VED, and the orientation is gradually transformed into < 111 > crystallographic direction families at a higher VED. Besides, the orientation relationship between A2 and B2 phases is {100}A2∥{100}B2. 5) The micro-hardness and the corrosion resistance are gradually improved as the VED rises and the maximum micro-hardness rises to 632.8HV, which is much higher than that of the samples fabricated by other methods (> 100Hv). The main reason is that the higher VED results in the much more rapid cooling rate and the formation of more B2 phases.
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