Selective laser melted equiatomic CoCrFeMnNi high-entropy alloy: Microstructure, anisotropic mechanical response, and multiple strengthening mechanism

Journal of Alloys and Compounds 805 (2019) 680e691

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Selective laser melted equiatomic CoCrFeMnNi high-entropy alloy: Microstructure, anisotropic mechanical response, and multiple strengthening mechanism Young-Kyun Kim a, Jungho Choe b, Kee-Ahn Lee a, * a b

Department of Materials Science and Engineering, Inha University, Incheon, 22212, Republic of Korea Korea Institute of Materials Science (KIMS), Changwon, 51508, Republic of Korea

a r t i c l e i n f o

a b s t r a c t

Article history: Received 7 May 2019 Received in revised form 9 July 2019 Accepted 11 July 2019 Available online 12 July 2019

One of the major challenges of equiatomic CoCrFeMnNi HEA is to manufacture parts with complex geometry that have higher yield strength. Equiatomic CoCrFeMnNi HEA was successfully fabricated in the present study with selective laser melting. The unique microstructure and mechanical anisotropy that generally appear in additive manufactured materials were investigated. SLM-built HEA has strongly oriented grains, dislocation networks, and nano-sized oxides. In addition, the average grain sizes were measured as 15.66 mm, 12.93 mm, and 5.98 mm on the plane perpendicular to the scanning direction (SD), transverse direction (TD), and building direction (BD), respectively. A compressive test measured outstanding yield strengths (YS) of 778.4 MPa, 766.4 MPa, and 703.5 MPa in the loading axis of SD, TD, and BD, respectively. These outstanding YSs are the result of a combination of fine grain sizes, high dislocation density and nano-sized oxides. In addition, anisotropy in mechanical properties are result from different values of Taylor factor and grain size according to the loading axis. After a compression test, the geometrically necessary dislocation density was found to differ about 2.5 times on each plane parallel to the loading axis in the same macro strain. Based on such findings, the relationship among microstructure, mechanical anisotropy and deformation mechanism are discussed in the present paper. Furthermore, the mechanical properties of SLM-built equiatomic CoCrFeMnNi HEA are predicted by using multiple strengthening mechanisms considering the microstructural characteristics. © 2019 Published by Elsevier B.V.

Keywords: Selective laser melting High-entropy alloy Microstructure Anisotropy Mechanical properties Strengthening mechanisms

1. Introduction High-entropy alloys (HEAs), also referred to as multi-principal element alloys (MPEAs) and compositionally concentrated alloys (CCAs), are highly regarded as being able to replace traditional metallic materials owing to their superior mechanical, physical and chemical properties [1e7]. Unlike conventional alloys, where alloying elements are added to the “principal element,” HEAs are unique materials composed of equiatomic or near-equiatomic mixing of multiple elements. Furthermore, HEAs have high mixing entropy in available temperature ranges and maintain a very stable random solid solution structure [8]. Interest in HEAs has increased significantly in recent years, and its thermal, physical and chemical properties have been investigated in many studies [9e14].

* Corresponding author. E-mail address: [email protected] (K.-A. Lee). https://doi.org/10.1016/j.jallcom.2019.07.106 0925-8388/© 2019 Published by Elsevier B.V.

The study fields related to HEAs are expanding to include nonequiatomic medium-entropy alloy [15], oxide dispersion strengthened HEA [16] and interstitial HEAs [17]. However, the HEAs commonly used today are manufactured by casting [18,19] and powder metallurgy [20,21]. As a result, there is a limit in manufacturing complex shapes and controlling physical properties, and material waste occurs during the manufacturing process. Also, it is difficult to manufacture homogeneous and defect-free HEAs with complex geometry using traditional manufacturing processes. Recently, there have been attempts to manufacture HEAs using additive manufacturing (AM), one of the key fields of the Fourth Industrial Revolution [22e26]. AM is an emerging technology used to manufacture near-net shape parts by layering powders in a layerby-layer incremental manner based on computer-aided design (CAD) models [27,28]. There are almost no limits to the shapes that can be manufactured with AM, allowing it to be considered as a method for manufacturing high-performance and customized parts for the aerospace, defense, automotive and biomedical fields.

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Among the known metal AM techniques, the most popular process is selective laser melting (SLM), which is a powder bed fusion method. Compared to direct energy deposition or electron beam melting, SLM is known to have a relatively faster cooling rate, which forms a unique microstructure, and to produce outstanding mechanical properties [29,30]. In addition, SLM has the advantage of greater dimensional accuracy and ability to control a uniform microstructure. If it is employed for manufacturing HEAs, additive manufacturing is expected to create a synergy effect, and recent studies on manufacturing equiatomic CoCrFeMnNi HEA using metal AM have been conducted [22e24,26]. Most of the studies conducted in relation to AM-built HEA have focused on material optimization and microstructural changes with process condition control. Chew et al. [22] manufactured equiatomic CoCrFeMnNi using laser aided AM and reported its microstructural and roomtemperature/cryogenic-temperature tensile properties. Zhu et al. [23] conducted a study on the influences of process parameters and heat treatment on the microstructure and tensile properties of SLM-built equiatomic CoCrFeMnNi HEA. Also, Li et al. [31] analyzed the influence of process parameters and hot isostatic pressing on changes in microstructural and tensile properties. The three studies all confirmed that AM-built HEA has epitaxial grain growth or a hierarchical microstructure due to directional solidification. This hierarchical microstructure can cause mechanical anisotropy, but no studies have been conducted in relation to this. In particular, in order to use SLM-built HEA as a structural material, it is necessary to identify the cause of mechanical anisotropy based on initial microstructural characteristics and analyze the difference in deformation behavior according to the loading axis. But as of now, no study has been conducted in relation to this. In the present study, equiatomic CoCrFeMnNi HEA was prepared using selective laser melting, and its compressive mechanical properties were examined along the three directions of an as-built sample to understand the mechanical anisotropy and its relationship to the unique microstructure. In addition, the mechanical properties of SLM-built equiatomic CoCrFeMnNi HEA were predicted by using multiple strengthening mechanisms. 2. Experimental 2.1. Material and SLM manufacturing conditions Fig. 1 shows the properties of the equiatomic CoCrFeMnNi HEA powder used for the SLM process. Powder with purity greater than 99.9 wt% was used, and its shape was spherical (Fig. 1a). The average particle diameter of the initial powder was measured as

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27.2 mm (d10 ¼ 17.9 mm, d50 ¼ 26.1 mm, and d90 ¼ 38.2 mm) (Fig. 1b). EBSD-EDS analysis results of the initial powder confirmed that the average grain size inside a powder particle was 3.2 mm, and all elements were evenly distributed (Fig. 1c). Equiatomic CoCrFeMnNi HEA powders were then used as feedstock for SLM processing (model: Concept Laser Mlab Cusing) under high-purity Ar gas to minimize oxidation. A laser scanning speed of 600 mm/s was applied to manufacture bar-type samples with dimensions of 85 mm  10 mm  10 mm. A laser power of 90 W with a hatch space of 80 mm was applied. The layer thickness was set to 25 mm. A scanning pattern (i.e., 180 angle rotation for the two consecutive layers) was also applied. Volumetric energy density (VED) is considered a key factor for the properties of SLM-built materials, and is defined as follows [31,32]:

VED ¼

P vht

(1)

where P is the laser power, v is the scan speed, h is the hatch spacing, and d denotes the layer thickness. The VED is measured to 75 J/mm3 in the present study. A macroscopic image of the SLMbuilt HEAs is shown in Fig. 2, and no cracks or defects were observed on the surface.

2.2. Microstructural characterization To identify the phases of SLM-built equiatomic CoCrFeMnNi, an X-ray diffractometer (XRD, Ultima IV, Cu Ka radiation, scan step size; 0.02 deg., scan rate; 2 deg. min1) was used. The dislocation density of as-built equiatomic CoCrFeMnNi was calculated using the convolutional multiple whole profile (CMWP) method [33]. To investigate the initial and deformed microstructures of SLM-built equiatomic CoCrFeMnNi alloy, samples were ground with #400 ~ #1200 grit silicon carbide papers and 1 mm diamond suspension, and then mirror polished with a colloidal silica (CS) suspension. Back-scattered electron imaging (BSEI) was performed in field emission scanning electron microscopy (FE-SEM, MYRA 3 XMH, Tescan, Czech Republic). In addition, crystallographic orientation and grain size was analyzed by electron backscatter diffraction (EBSD, Nordlys-CMOS detector, binning: 2  2, Oxford, United Kingdom). Dislocation structures were analyzed by electron channeling contrast (ECC) imaging at an acceleration voltage of 30 kV with a BSE detector. The working distance of ECC imaging was set to 7 mm with a sample tilt of 2.6 . To observe nano-sized phases in SLM-built equiatomic CoCrFeMnNi, a Cs-corrected scanning transmission electron microscope (STEM, JEM-ARM200F, Japan) and field emission TEM (FE-TEM, JEM-2100F, Japan) were used.

Fig. 1. Raw equiatomic CoCrFeMnNi HEA powders: (a) SEM morphology, (b) distribution of powder size, and (c) EBSD IPF map of initial powder.

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Fig. 2. (a) Schematic illustration of the scanning strategy. BD, SD, and TD indicate building direction, scanning direction and transverse direction, respectively. (b) Macroscopic image of as-built sample.

2.3. Compression testing To evaluate the room-temperature compressive properties of SLM-built equiatomic CoCrFeMnNi, a cylindrical sample with a size of F 4-mm  6-mm height was electrical discharge machined. Room-temperature compression tests were performed using MTS810 equipment with different directions of scanning direction (SD), transverse direction (TD), and building direction (BD). An initial strain rate of 1  103 s1 and an engineering strain limit of 0.4 were applied (i.e., compression testing was implemented until the strain reached 40%). The compression test was performed three times in each condition, and the most representative curve was used. Here, MoS2 spray was used as a lubricant to reduce physical friction.

3. Results 3.1. Microstructure of additively manufactured equiatomic CoCrFeMnNi high-entropy alloy 3.1.1. Phase and elemental composition analysis results Fig. 3 shows the X-ray diffraction results and fitted values of fullwidth at half-maximum (FWHM) of selective laser melted equiatomic CoCrFeMnNi HEA. First, the XRD analysis only identified a single FCC phase, and no other phases were detected (Fig. 3a). In

addition, FWHM calculated from each diffraction plane of additively manufactured HEA had a relatively higher value compared to conventionally manufactured HEA [34], which indicated the possibility of it having more lattice defects. Dislocation density measured using CMWP was 1.88  1014m2, which is approximately two times greater than that of conventionally manufactured HEA, hot rolled, and homogenized equiatomic CoCrFeMnNi HEA (6.34  1012m2) [34]. In general, the additive manufactured sample had high dislocation density, and this is reportedly formed by the fast cooling and solidification rates [35]. Therefore, relatively higher strength can be induced in the additive manufactured sample by having higher dislocation density than commercial materials [36]. Fig. 4 shows EDS mapping and an EBSD phase map. First, EDS mapping (Fig. 4a) identified an even distribution of Co, Cr, Fe, Mn and Ni elements within the material, and all elements were measured near 20 at.%. This means that the elemental segregation that is reported to occur in some SLM-built samples did not occur in this study. The phase map of a 600 mm  1200 mm plane perpendicular to BD was analyzed using EBSD, and as with the XRD analysis (Fig. 3a), only a single FCC structure was identified and no additional phases were identified. Based on the findings from EBSD, XRD and EDS analysis, SLM-built equiatomic CoCrFeMnNi HEA was confirmed to have a single FCC structure and compositional homogeneity.

Fig. 3. (a) X-ray diffraction patterns and (b) full-width of half-maximum (FWHM) of SLM-processed equiatomic CoCrFeMnNi HEA.

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Fig. 4. (a) SEM-BSE image and elemental distribution maps of the area and (b) EBSD phase map.

3.1.2. SLM microstructure and dislocation network Fig. 5 represents typical EBSD inverse pole figure (IPF) maps in the SLM-built equiatomic CoCrFeMnNi with different planes perpendicular to the building direction (BD, Fig. 5a), scanning direction (SD, Fig. 5b), and transverse direction (TD, Fig. 5c),

respectively. In the EBSD IPF maps of the BD and SD planes (Fig. 5a and b), an aggregation of <110> and <001> orientations was dominant parallel to the building direction. These grains grow epitaxial to the building direction and form a hierarchical microstructure. In contrast, the EBSD IPF map of the TD plane (Fig. 5c)

Fig. 5. EBSD IPF maps of (a) BD plane, (b) SD plane, and (c) TD plane in as-built equiatomic CoCrFeMnNi HEA.

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shows <001> and <110> orientations, while also revealing a large amount of <111> orientation grains parallel to the building direction. In general, additive manufactured samples are known to grow easily in the preferred growth direction of <001> for the FCC structure in the direction of the solidification outgrow when powder is being melted and solidified [37]. However, in cases where the scanning strategy is controlled to a degree, as in this study, it is confirmed that the hierarchical microstructure is identical, but the preferred orientation can be controlled to some degree (Fig. 5). Also, the average grain size was different for each direction, measuring 5.98 mm, 15.66 mm, and 12.93 mm, respectively, for the BD, SD and TD planes. While each plane does have areas with massive grain sizes, since the ratio of areas with grains smaller than 10 mm was high, the average grain size was relatively low (Fig. S1). Fig. 6 shows a high-magnification BSE image of the SLM-built HEA. BSE image analysis identified cellular- and columnar-type structures within the grain. In general, additive manufactured materials are known to have a unique cell structure, and the cell size is known to vary according to the process condition [38]. According to recent studies by AlMangour et al. [39] and Wang et al. [40], the shape of the cell structure is formed due to either elemental segregation or dislocation network. Fig. 7 shows typical cellular- and columnar-type ECC images in SLM-built equiatomic CoCrFeMnNi. As in the recent studies, the cellular structure was confirmed to have been formed by a dislocation network (Fig. 7a). The cell structure measured approximately 413.2 ± 79.5 nm, and some dislocations were observed within the cell. In addition, there was a fine phase with an average size of 27.3 ± 3.2 nm in the cell structure boundary and inner cell structure area. Such small sizes were the reason why the XRD and EBSD analysis results did not identify additional phases. In the ECC image of the columnar structure (Fig. 7b), it is clear that a dislocation network caused a cellular- or columnar-structure boundary. In other words, the reason why the material has greater dislocation density compared to homogenized HEA is the dislocation network formed by rapid solidification during the SLM process. The average grain size of the BD, SD, and TD planes were different, but the size and shape of cellular and columnar structures were very similar in all three phases.

3.1.3. TEM analysis results Fig. 8 represents the combined STEM and EDS analysis results of

Fig. 6. High magnification BSE image showing the cellular and columnar structures in SLM-processed equiatomic CoCrFeMnNi HEA.

the SLM-built HEA. As seen in the ECC image, phases of a few tens of nm were observed, and they were confirmed to be Mn-rich oxide via EDS mapping. In general, equiatomic CoCrFeMnNi manufactured with a casting process reportedly forms CreMn-based oxides with a few mm [11]. Studies have also reported that materials heat treated at a particular temperature form s-phase [41]. However, no study has reported that additively manufactured CoCrFeMnNi HEA forms a few tens of nm-sized Mn-rich oxide. No elemental segregation was observed in the cellular structure boundary mentioned in section 3.1.2, and this suggested the possibility that the cellular structure can only be formed by dislocation networks. Fig. 9 shows HR-TEM and corresponding fast Fourier transform (FFT) and inverse FFT (IFFT) results of Mn-rich oxide in additively manufactured equiatomic CoCrFeMnNi HEA. HR-TEM image analysis identified oxides approximately 19.8 nm in size, and Mn2O3 phase was confirmed through FFT and IFFT. At the same time, (431) and (222) planes were analyzed, and the d-spacing of each plane was 0.184 nm and 0.272 nm, respectively. In other words, an evenly distributed phase with high density was identified as Mn2O3 through EDS mapping and HR-TEM. 3.2. Mechanical properties and work hardening behavior Fig. 10 shows the representative compressive stress-strain curves of the SLM-built HEAs with different loading directions. The SLM-built HEA had similar yield strength values in the loading axes of both SD and TD (SD: 778.4 MPa and TD: 766.4 MPa), while the loading axis of BD had ~75 MPa lower yield strength than that of SD compression (Fig. 10a). In addition, the stress-strain curves were also similar in TD and SD compression, but different in BD compression. In other words, additively manufactured equiatomic CoCrFeMnNi HEA was confirmed to form mechanical anisotropy. The yield strength of materials manufactured in this study was compared with the yield strength of additively manufactured HEAs [22,26,31] and annealed HEA [42] reported up to the present, and the results are shown in Fig. 10b. In the case of materials manufactured with SLM [31], 519 MPa for the as-built state and 485 MPa for as-built þ HIP were reported. Materials manufactured with laser aided additive manufacturing [22] and laser metal deposition [26] were confirmed to have yield strength of 518 MPa and 290 MPa, respectively. Unlike this, the sample manufactured with the conventional process and annealed was confirmed to have a yield strength of 254 MPa [42]. In other words, the material manufactured in this study was confirmed to have relatively higher strength compared to other additively manufactured and conventionally manufactured equiatomic CoCrFeMnNi HEAs. The difference in such mechanical properties is suspected to be the result of microstructural characteristics or differences in defects caused by different manufacturing processes. Related strengthening mechanisms are discussed in section 4.2. In order to achieve reliability of the compressive yield strength used, tensile tests were performed on identical materials, and the results are shown in Fig. S2. It was confirmed that there was no significant difference in yield strength under tensile and compressive loads, and this finding allowed the comparison of mechanical properties using compressive yield strength. Fig. 11 shows true stress-strain curves (Fig. 11a), work hardening rate (WHR, ds/dε) curves (Fig. 11b), and corresponding surface observation results after compression at a strain up to 0.4. Upon reviewing the true stress-strain curve, SD and TD compressions with similar mechanical properties showed similar deformation behaviors (Fig. 11a). In the case of BD compression, the yield strength was relatively low, but the flow stress increase in the plasticity region is large, and peak stress was also confirmed to be the highest. Upon reviewing the WHR curves, the SD, TD and BD

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Fig. 7. ECC images showing the dislocation networks and fine oxides in SLM-processed equiatomic CoCrFeMnNi HEA: (a) cellular structure and (b) columnar structure.

Fig. 8. STEM image and elemental distribution maps of STEM image area showing fine Mn-based oxide.

Fig. 9. (a) An HRTEM image showing the nano-sized Mn-based oxide and (b, c) FFT pattern and Fourier filtered image of the selected square region in (a).

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Fig. 10. (a) Typical compressive stress-strain curves of the SLM-processed equiatomic CoCrFeMnNi HEAs with different loading axes. (b) The engineering yield strength data in this study and other additively manufactured equiatomic CoCrFeMnNi from the literature. SLM refers to selective laser melting, HIP is hot isostatic pressing, LAAM is laser aided additive manufacturing, LMD is laser metal deposition and anneal. is annealing.

Fig. 11. (a) True stress-strain curves and (b) work hardening rate curves of SLM-processed equiatomic CoCrFeMnNi HEAs. Surface observation results after compression test with different loading axes: (a) SD, (b) TD, and (c) BD.

had significant differences like those found in the true stress-strain curves (Fig. 11b). SD compression, which showed the highest yield strength, had the lowest WHR, and BD compression, which had the lowest yield strength, had a relatively higher WHR. In other words, the WHR had an opposite tendency to the yield strength. In the initial deformation stage ðεt ¼ 0 0:26Þ in SD compression, the WHR rapidly decreased as plastic strain increased. On the other hand, in the deformation stage (εt > 0:26Þ, the WHR tended to increase gradually as deformation progressed. TD compression also showed two deformation stages during the compression test, but represents a relatively narrow stage II. In contrast, in BD compression, the WHR tended to decrease monotonically as

deformation progressed. In general, the WHR reportedly changes according to the deformation mechanism change of equiatomic CoCrFeMnNi HEA during deformation (i.e., creation of deformation twinning) [19,43]. It is possible to predict that deformation twins (DTs) capable of influencing deformation behavior in SD and TD compressions will be formed. Based on such findings, SLM-built HEA was confirmed to not only have a significant change in yield strength, but also flow stress and WHR due to differences in microstructural characteristics along directions. To analyze the characteristics of mechanical anisotropy, the surfaces after compressive deformation were observed, and the results are shown in Fig. 11c and d. Surface observation results

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show the geometric differences in compressed samples with different loading directions. SD and TD compressed samples, which have a relatively lower WHR during deformation, deformed severely in a particular direction, while BD compression, which has a higher WHR, deformed evenly. The reason for this is suspected to be due to it being easier for a particular plane to accept deformation while a compressive load is being applied. The reason why deformation progressed in an even mode in BD compression is thought to be the result of similar microstructures (grain size, Taylor factor) of the two planes (TD and SD) placed perpendicularly. 3.3. Microstructural evolutions during compressive deformation Fig. 12 shows a geometrically necessary dislocations (GNDs) distribution map (Fig. 12a and b) and GNDs density (Fig. 12c) of the SD compressed sample. Here, GNDs were calculated by using Nye's dislocation tensor [44,45]. EBSD analysis performed on the representative two planes confirmed that the plane with relatively more strain accumulation during SD compression was the TD plane, and the plane with less deformation was the BD plane. In the case of the TD plane, dislocations concentrated along the high angle grain boundaries (HAGBs, black line), and sub-grains were observed within the grain. In the case of the BD plane, while the dislocations concentrated on the grain were identical, no sub-grains were formed and dislocation density increased at particular grains. The grains where dislocations were concentrated were identified as fine grains within the track boundary. The GNDs density measured 1.48  1014 m2 in the as-built sample and 4.55  1014 m2 and 1.18  1015 m 2 in the BD and TD planes of the compressed sample. In other words, the evolution of GNDs was confirmed to be relatively larger in the TD plane, and therefore it is easier for the TD plane to accommodate deformation (strain accumulation) than the BD plane, and this results in a unique deformation behavior according to direction (uneven shape of the compressive sample). Fig. 13 shows high-magnification EBSD band contrast, IPF, and GND distribution maps of the planes perpendicular to TD (Fig. 13aec) and BD (Fig. 13def) of the SD compressed sample. Band

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contrast map analysis of the TD plane identified additional HAGBs within the grain (Fig. 13a). IPF map observation confirmed that such HAGBs are S3 DTs that have a q z 60 relationship to neighboring grains (Fig. 13b). In addition, the GNDs distribution map identified dislocation evolution within DTs along with the initial grain (Fig. 13c). The reason why high dislocation density is found within the interior of DTs is known to be a result of the large amount of plastic deformation accommodated after the formation of DTs. Band contrast map analysis of the BD plane did not identify additional HAGBs (Fig. 13d). The IPF map (Fig. 13e) also did not identify additional grains aside from the initial grain, and a relatively smaller dislocation density increase was confirmed with the GNDs map (Fig. 13f). Based on such findings, it is suspected that strain accumulation during SD compression occurs more easily in the TD plane, and while the total compressive strain is identical, the true applied strain is predicted to be higher than the BD plane. This characteristic can be found from the difference between dislocation densities, and as stored strain energy is higher in the TD plane, this is thought to cause the difference in DTs formation. Fig. 14 shows typical ECC images of the TD plane after the SD compression test. Unlike the findings from EBSD, the ECC image did not identify massive DTs. In the case of the ECC image, individual fine DTs created bundles that formed a band shape. As a result, EBSD (Fig. 13) recognized multiple DTs as a single DT because of the difference in step size. The thickness of these individual DTs was measured as 12.6 ± 2.3 nm in the high magnification ECC image. Also, the high dislocation density in the EBSD GNDs distribution map (Fig. 13) can be explained with the multiple dislocations formed between DTs identified in the ECC results (Fig. 14). In the case of the cellular and columnar structures found in the initial microstructure, they had a tendency to disappear after compressive deformation, and this is suspected to be due to the failure to maintain the cellular structure during dislocation accumulation. If deformation is accommodated beyond a particular strain, a driving force to create a sub-grain of more than 1 mm, like that found in the EBSD GNDs distribution map (Fig. 13c) is formed, and this driving force is thought to be the cause of the destruction of the cellular structure.

Fig. 12. Geometrically necessary dislocation distribution at different planes of the SD compression sample: (a) TD plane and (b) BD plane; (c) calculated GNDs density.

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Fig. 13. EBSD band contrast maps (a, d), IPF maps (b, e) and GNDs distribution maps (c, f) showing the deformed microstructures of SD compression.

Fig. 14. ECC images showing bundles of deformation twins in the TD plane after SD compressive deformation.

4. Discussion 4.1. Effect of Taylor factor on mechanical anisotropy SLM-built equiatomic CoCrFeMnNi HEA was confirmed to have different average grain sizes according to its direction (Fig. 5aec). However, there is no difference in cellular and columnar structure according to direction, and the dislocation density and fine oxides in the tens of nm were confirmed to be similar (Figs. 3, 7 and 8). In addition, elemental segregation (i.e., segregation engineering [39,46], which is considered to be a strengthening mechanism at present, did not occur. Thus, we have focused on the Taylor factor according to the loading axis of SLM-built HEA. The Taylor factors of additively manufactured CoCrFeMnNi HEA calculated by AZTec HKL are shown in Table 1. Fig. 15 is the analysis results of the Taylor factor of the planes perpendicular to BD and TD for SD compression, which has the

Table 1 Taylor factor of the planes perpendicular to scanning, transverse, and building directions at different loading axes.

Loading axis: SD Loading axis: TD Loading axis: BD

BD plane

TD plane

SD plane

3.328 3.340 e

3.113 e 3.012

e 3.143 3.050

highest yield strength. The BD plane (3.328, Fig. 15a) and TD plane (3.113, Fig. 15b) had very different Taylor factors, and the Taylor factor distribution map (Fig. 15c) shows the differences in their distribution. As shown in Table 1, the Taylor factor changes significantly as the loading axis changes in the BD, TD and SD planes. In previous reports, equiatomic CoCrFeMnNi HEA had a Taylor factor of 3.06 [44]. However, the BD plane in this study had a Taylor factor of 3.328 and 3.340 in SD and TD compressions, respectively, which was a significantly different result from other reports. In general, if <111>-fiber texture and <100>-fiber texture are present in FCC polycrystals, the Taylor factors are 3.67 and 2.45, respectively [47]. In other words, the difference in Taylor factors is thought to be the critical cause of mechanical anisotropy according to the loading axis, and this is explained by texture strengthening. Therefore, the differences in mechanical anisotropy and deformation behavior can be explained by the different Taylor factors in planes of SLM-built HEA according to each loading axis. Furthermore, the Taylor factor can also influence the critical stress for twinning. According to Laplanche et al. [43], equiatomic CoCrFeMnNi HEA has critical stress for twinning that is reported as 720±30 MPa, and many other researchers have also measured and presented twinning stress [48,49]. In the case of equiatomic CoCrFeMnNi HEA, critical stress for twinning reportedly differs according to grain size [50,51], and recently, Sun et al. [52] presented a prediction model considering grain size as follows:

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Fig. 15. Taylor factor maps for loading along SD: (a) BD plane, (b) TD plane and (c) corresponding Taylor factor distribution map in (a) and (b).

sT ¼ M

g bp

k þ pTffiffiffi d

(2)

where sT is the critical stress for twinning stress, g is stacking fault energy (21 mJm2 [53]), kT represents the Hall-Petch constant for pffiffiffiffiffiffiffi twinning (980 MPa, mm [54]), bp denotes the Burgers vector of a partial dislocation (1.46  1010 m [55]), d is average grain size (BD: 5.98 mm, SD: 15.66 mm, and TD: 12.93 mm), and M is the Taylor factor (See Taylor factors in Table 1). sT changes according to Taylor factor, and in the case of the BD plane, its value is predicted to be ~880 MPa in SD compression. In the case of the TD plane, which has an sT of ~719 MPa, it is suspected that DTs evolution will have greater deformation (Fig. 13). However, the calculated sT of SLM-built HEA was significantly different from the actual WHR increase. This is thought to be due to the complex interaction of strain accumulation inhomogeneity and unique initial microstructure (e.g., dislocation network, fine oxides, and different grain sizes along the direction). 4.2. Multiple strengthening mechanisms Until now, no studies have reported the strengthening mechanism of additively manufactured equiatomic CoCrFeMnNi HEA other than a study by Zhu et al. [23]. Their study explained the three mechanisms of lattice friction stress (sf ), grain boundary strengthening related to the Hall-Petch relationship (sGB ) and dislocation strengthening (sr ) as contributing to strengthening. However, this yield strength prediction model was later discovered to have limitations in making accurate predictions. The reason for these limitations is thought to be the unique initial microstructure SLM-built HEA has compared to cast or wrought HEAs. In addition, as dislocation density is measured using the cell size-dislocation density formula presented by Kocks-Mecking, this is also suspected to have contributed to the difference [56]:

pffiffiffi

r ¼ c=l

sy ¼ sf þ sSS þ sGB þ sOr þ sr

(4)

First, sf of equiatomic CoCrFeMnNi is reported as 125 MPa in the literature [57]. In addition, Moon et al. [58] reported that intrinsic yield strength is 147 MPa in equiatomic CoCrFeMnNi HEA. Thus, the estimated contribution of solid solution strength is 22 MPa. The average grain size significantly affects the yield strength of a metallic material because HAGBs suppress dislocation motion. Therefore, grain boundary strengthening can be calculated by the Hall-Petch relation, sGB , which contributes to the overall strengthening and is expressed as [59]:

sGB ¼ kd1=2

(5)

where k denotes the strengthening coefficient and d is the average grain size of the SLM-built HEA. For the SLM-built HEA, the strengthening coefficient is ~494 MPa mm1/2 [57] and the average grain sizes are measured as 5.98 mm, 15.66 mm, and 12.93 mm in the BD, SD, and TD planes, respectively. The estimated strengths of the BD, SD, and TD planes of the as-built sample from the grain boundary strengthening contribution using equation (5) are 202.1 MPa, 124.8 MPa, and 137.4 MPa, respectively. Another strengthening mechanism for the superior strength of the SLM-built equiatomic CoCrFeMnNi HEA is Orowan strengthening via nano-sized oxides in the matrix, given by Ref. [39]:

0:4M Gb sOr ¼ pffiffiffiffiffiffiffiffiffiffiffi, ,ln p 1v L

qffiffiffi 2d 3 p b

(6)

Where L is the interparticle spacing and is calculated using equation (7), given below:

L¼

rffiffiffi rffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2 p dp ð  1Þ 3 4vp

(7)

(3)

where r is dislocation density, c is a constant and l represents cell size. However, in order to predict the actual yield strength of SLMbuilt HEA, solid solution strengthening (sSS ) and Orowan strengthening related to oxides (sOr ) must also be considered along with the strengthening contributions above. Thus, we have considered the five strengthening contributions for the additively manufactured equiatomic CoCrFeMnNi HEA given as below:

and M is the Taylor factor at each plane of SLM-built HEA, and is summarized in Table 1. v denotes Poisson's ratio (0.26 [60]), b is the Burgers vector (0.255 [43]), dp represents the average particle diameter (27.3 nm), G is the shear modulus (80.85 GPa calculated by 85-(e448/T-1) [61]), and vp is the volume fraction of the nano-sized oxides (0.054). The estimated strengths by Orowan strengthening in the SLM-built HEA vary along the plane and loading axis. Therefore, we have focused on the SD and BD compressions, which

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show the biggest differences in strength, and the estimated strengths of the BD and TD planes are 554.6 MPa and 518.8 MPa in SD compression. In the case of BD compression, the estimated strengths are 508.3 MPa and 501.9 MPa at the SD and TD planes. The final strengthening mechanism in the SLM-built HEA is dislocation strengthening, sr , given by Ref. [43]:

pffiffiffi

sr ¼ MaGb r

(8)

where a is the constant (0.33 [42]) and r is the dislocation density (1.88  1014 m2 by CMWP method). Using equation (8), it is confirmed that the contribution of dislocation strengthening to the strength of the SLM-built HEA also varies along the loading axis at each plane. Researchers have focused on the SD and BD compressions, as mentioned above, and the estimated strengths of the BD and TD planes in SD compression are 281.1 MPa and 262.9 MPa, respectively. In contrast, the estimated strengths of the SD and TD planes are 257.6 MPa and 254.4 MPa in BD compression. In other words, it was also confirmed that the strength contribution is different because it has a different Taylor factor on the same plane. Using equation (4), the overall strength of the SLM-built equiatomic CoCrFeMnNi HEA is somewhat higher than experimentally measured strengths. This phenomenon is due to the fact that grain boundary, Orowan, and dislocation strengthening mechanisms interact when an external force is applied. Thus, many researchers regard the root mean square of these effects as below [39]:

sy ¼ sf þ sSS þ

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi sGB 2 þ sOr 2 þ sr 2

(9)

The total strength is calculated based on the contributions of all five strengthening mechanisms for SLM-built HEA. Since the contribution in each plane differs depending on the loading axis, the strength is predicted using a mean value such as sx ¼ sxðPlane 1Þ þsxðPlane 2Þ . Here, sx are the strengthening mechanisms (i.e., 2 grain boundary strengthening, Orowan strengthening, and dislocation strengthening). For SD compression, the estimated total strength is 772.2 MPa, which is similar to the experimentally measured strength of 778.4 MPa. In the case of BD compression, the estimated total strength is 728.2 MPa, which is also similar to the experimentally measured strength of 703.5 MPa. The contributions of each strengthening mechanism to the overall predicted strengths are shown in Fig. 16. These predicted strengths are similarly better with the experimentally measured values than the simply added strength and previous yield strength prediction study [23]. In other words, yield strength can be predicted considering various microstructural characteristics in SLM-built HEA. It was also confirmed that the contribution of Orowan strengthening is the highest at ~44%, indicating that an oxide plays a key role in the superior mechanical properties of additively manufactured equiatomic CoCrFeMnNi HEA. 5. Conclusions The correlations between microstructure and mechanical anisotropy of selective laser melted equiatomic CoCrFeMnNi highentropy alloy were investigated. The following conclusions were drawn from the results of this study: 1. SLM-built equiatomic CoCrFeMnNi HEA has a single FCC structure and state of compositional homogeneity in the as-built sample. The average grain size was measured as 5.98 mm, 15.66 mm, and 12.93 mm in the BD, SD, and TD planes, respectively. In addition, cellular and columnar structures were formed by dislocation network. Oxides with an average size of 27.3 nm were also formed in SLM-built HEA.

Fig. 16. Comparison of the contributions of each strengthening mechanism to the SD and BD compression of SLM-processed equiatomic CoCrFeMnNi HEA. SD represents the plane perpendicular to scanning direction and BD denotes the plane perpendicular to building direction.

2. The yield strengths of SLM-built HEA in this study represent higher values than other additively manufactured equiatomic CoCrFeMnNi HEA. In particular, a yield strength of 778.4 MPa can be achieved in scanning direction compression. It was also confirmed that the yield strength has an anisotropic response along each loading axis (BD, SD, and TD). 3. For SD and TD compressions, deformation of the BD plane is accommodated by dislocation slip, whereas deformation twinning occurs at the plane perpendicular to the BD plane. This phenomenon is associated with the different grain sizes and Taylor factors at each loading axis (i.e., critical stress for twinning affected by grain size and Taylor factor). In the case of SD compression, the BD plane and TD plane have a Taylor factor of 3.328 and 3.113, respectively. 4. This study predicted yield strength considering the Taylor factor, nano-sized oxides and grain size, which were not considered in previous SLM-built HEAs strengthening models. The newly proposed yield strength prediction model is more comprehensive than the general model for predicting the yield strength of additively manufactured equiatomic CoCrFeMnNi HEAs, and these results are in good agreement with the experimentally measured yield strength. Acknowledgement This study was supported by the Fundamental Research Program of the Korea Institute of Materials Science (Grant No. PNK5520). It was also supported by a National Research Foundation of Korea (NRF) grant funded by the Korean government (MEST) (No. 2019R1A2C1008904). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jallcom.2019.07.106. References [1] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448e511. [2] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.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) 299e303. [3] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development

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