Fabricating Homogeneous FeCoCrNi High-Entropy Alloys via SLM In Situ Alloying

: Selective laser melting (SLM) in situ alloying is an effective way to design and fabricate novel materials in which the elemental powder is adopted as the raw material and micro-areas of elemental powder blend are alloyed synchronously in the forming process of selective laser melting (SLM). The pre-alloying process of preparation of raw material powder can be left out, and a batch of bulk samples can be prepared via the technology combined with quantitative powder mixing and feeding. The technique can be applied to high-throughput sample preparation to efﬁciently obtain a microstructure and performance data for material design. In the present work, bulk equiatomic FeCoCrNi high-entropy alloys with different processing parameters were fabricated via laser in situ alloying. Finite element simulation and CALPHAD calculation were used to determine the appropriate SLM and post-heating parameters. SEM (scanning electron microscope), EDS (energy dispersive spectroscopy), XRD (X-ray diffraction), and mechanical testing were used to characterize the composition, microstructure, and mechanical properties of as-printed and post-heat-treated samples. The experimental results show that the composition deviation of laser in situ alloying samples could be controlled within 20 wt %. The crystal structure of as-printed samples is a single-phase face-centered cubic (FCC), which is the same as those prepared by the traditional method. The mechanical properties of the samples prepared by laser in situ alloying with elemental powder blend are comparable to those prepared by pre-alloying powder and much higher than those prepared by the traditional method (arc melting). As-printed samples can get a homogeneous microstructure under the optimal laser in situ alloying process combined with post-heat treatment at 1200 ◦ C for 20 h.


Introduction
Most structural metallic materials are based on a primary element to improve the overall performance by mixing the primary element with other elements. The pioneering work of Cantor and Yeh et al. [1,2] on mixing multiple high concentration elements has opened up a new field in material science called high-entropy alloys (HEAs). HEAs usually contain five or more major elements with a concentration between 5 and 35 at %, which have superior high temperature strength, wear resistance, and oxidation resistance [3]. Representatives are the Co-Cr-Fe-Ni-Cu system with an FCC structure and the Al-Co-Cr-Fe-Ni system with a BCC (Body-centered cubic) structure. Traditional HEAs with an FCC structure usually have high plasticity, while the one with a BCC structure usually has high strength [4]. The definition of high-entropy alloys is proposed to have at least five main elements, but sometimes ternary CoCrNi alloy as well as quaternary CoCrFeNi homogeneously with the laser energy of 259 J/mm 3 . One of the core issues of the laser in situ alloying technique is the composition homogenization of samples during rapid melting and solidification, which directly affects the microstructure and final mechanical properties of the sample. In our previous study [20], the influence law of the laser in situ alloying process and heat treatment on the composition homogenization of Fe-Ni binary alloys was calculated by the CALPHAD method. Four groups of samples with different compositions were fabricated by SLM in situ alloying, and the average density was up to 99.8%. After high temperature heat treatment at 1200 • C or 1400 • C, the laser melting pool in the samples disappeared, grains with obvious boundaries presented in the samples, and the elements were distributed basically uniformly in the grain, which verified the feasibility of SLM in situ alloying technology.
In the present work, we fabricated quaternary equiatomic FeCoCrNi samples with different processing parameters via elemental powder blend and pre-alloyed powder, respectively. The size of molten pools was calculated by the finite element analysis method to select suitable laser in situ alloying parameters. The CALPHAD method with the TCHEA4 database were used to analyze the influence of element composition deviation on the matrix phase transformation. The modified SCHEIL model coupled with Thermo-Calc software was used to calculate the micro-segregation during the rapid solidification process. The microstructure and mechanical properties of samples synthesized by two kinds of raw material powders were compared by SEM, EDS, XRD, and a room temperature tensile experiment. To get homogeneous samples, DICTRA software, (Diffusion Module (DICTRA) Stockholm, Sweden) was used to calculate a reasonable high-temperature heat treatment schedule. The corresponding heat treatment experiments were carried out, and EDS analysis was used to verify the effect of heat treatment on element distribution. The present work provides a new method to fabricate bulk HEAs with homogeneous composition and can be applied to high-throughput synthesis technology.

Specimens Fabrication
The elemental powders and pre-alloyed powders used in the experiment were spherical ones prepared by vacuum atomization in which the purities were between 99.5% and 99.9%. The substrate was made of 304 L stainless steel with a thickness of 2 mm. Four kinds of elemental powders were mixed with an equiatomic ratio of 1:1:1:1 (converting to mass ratio is 25:26:23:26). To keep the powder free from contamination, a special plastic powder mixing barrel with built-in blades was printed. The barrel with powder blend was mixed at a high speed for 1.5 h to keep it fully mixed. The target element composition, elemental powder weight, and basic physical properties of the raw materials calculated by CALPHAD are shown in Table 1. The sphericity and particle size distribution of the elemental powders ( Figure 1) were determined by Prox-SE scanner (Phenomworld, Eindhoven, The Netherlands) and the BetterSize2000 laser particle size analyzer (Dandong Better Co., Ltd., Dandong, China).  SLM-DLM280 equipment produced by Hangzhou DeDi Intelligent Technology Co., Ltd. (Hangzhou, China) was adopted to print samples. The thickness of powder layer was set to 20 μm, the scanning line spacing was 80 μm, the laser spot diameter was 80 μm, and the scanning path adopted the reverse scanning strategy, in which the scanning path between layers is perpendicular. The laser molten pool shape and size under different printing process parameters were calculated with ANSYS additive software (2020R2, Ansys Inc., Pittsburgh, PA, USA) to explore a better processing parameter scope. The designed laser power varied from 100 to 250 W, while the scanning speed varied from 700 to 1100 mm/s. The final printing process parameters are shown in Table 2, in which elemental powder samples were indicated as HEE and pre-alloyed powder samples were indicated as HEA. In the process of SLM, powders were melted with extremely high temperature and then solidified rapidly (10 4 -10 5 °C/s) [21]; the microstructure and forming property of the samples were under the influence of various parameters, mainly including laser power (P, W), layer thickness (h, mm), scanning velocity (v, mm/s), and hatch spacing (s, mm), etc. For comprehensive evaluation of the forming property and comparability with reported data, volume laser energy density (VE = P/v·s·h, J/mm 3 ) was adopted as the evaluation criteria of laser input in this article.  SLM-DLM280 equipment produced by Hangzhou DeDi Intelligent Technology Co., Ltd. (Hangzhou, China) was adopted to print samples. The thickness of powder layer was set to 20 µm, the scanning line spacing was 80 µm, the laser spot diameter was 80 µm, and the scanning path adopted the reverse scanning strategy, in which the scanning path between layers is perpendicular. The laser molten pool shape and size under different printing process parameters were calculated with ANSYS additive software (2020R2, Ansys Inc., Pittsburgh, PA, USA) to explore a better processing parameter scope. The designed laser power varied from 100 to 250 W, while the scanning speed varied from 700 to 1100 mm/s. The final printing process parameters are shown in Table 2, in which elemental powder samples were indicated as HEE and pre-alloyed powder samples were indicated as HEA. In the process of SLM, powders were melted with extremely high temperature and then solidified rapidly (10 4 -10 5 • C/s) [21]; the microstructure and forming property of the samples were under the influence of various parameters, mainly including laser power (P, W), layer thickness (h, mm), scanning velocity (v, mm/s), and hatch spacing (s, mm), etc. For comprehensive evaluation of the forming property and comparability with reported data, volume laser energy density (VE = P/v·s·h, J/mm 3 ) was adopted as the evaluation criteria of laser input in this article.

Testing Methods
As-printed samples were separated from the substrate by linear cutting, and the density after cleaning and drying the samples was measured by Archimedes' drainage method. Each sample was measured two times, and the average value was taken. The macrostructure was observed by an Olympus X53 optical microscope (Olympus China Co., Ltd., Shanghai, China). The microstructure and morphology were investigated with a scanning electron microscope of Phenom ProX-SE desktop (Phenomworld, Eindhoven, the Netherlands). The composition mapping of the as-printed and heat-treatment samples was characterized by EDS simultaneously. Phase analysis of as-printed specimens was performed with a Bruker D8 Advance X-ray Diffractometer (Bruker Corporation, Billerica, Massachusetts, USA) using Co−K α (wavelength λ = 1.7902 Å). Hardness was measured using a VH-5 Vickers Hardness indenter (Everyone Instrument Co. Ltd., Shanghai, China) with a 500 g load for 10 s. The average hardness of each specimen was evaluated from 12 to 16 measurements. Specimens of non-standard sheets were used for tensile testing, as shown in Figure 3. To make the forming at one time, the specimens were printed with a size of 2 mm × 7 mm × 80 mm. The tensile testing temperature is 20 • C, and the elongation rate is 0.05 mm/s.

Thermodynamic Calculations
Thermo-Calc software (2021, Thermo-Calc Software AB, Stockholm, Sweden) and its associated TCHEA4 database were used for phase diagram calculation, the matrix phase structure and stable temperature range were analyzed, and the calculated phase structure was compared with XRD characterization. The Scheil-Gulliver solidification model is a commonly used computational model to evaluate microstructure segregation during the rapid solidification process [21]. The model assumes that the components in the system are homogeneously mixed in the liquid and there is no diffusion in the solid phase [22]. Those conditions do not exist in nature, so the result value is a theoretical limit. In the process of actual printing, mutual diffusion of the liquid phase and solid phase will lead to a redistribution of elements, which is expected to be less than the predicted segregation value of the Scheil-Gulliver model. This work adopted a revised Scheil back diffusion model embedded in Thermo-Calc software. The model considered the diffusion of interstitial atoms between liquid and solid phase and mutual diffusion among solid phases [23]. A more accurate prediction of alloying elements segregation in the solidification process of SLM can be worked out. The segregation results were coupled with DICTRA software to serve as initial composition before homogenization heat treatment. The distance of the DICTRA diffusion model was set to 100 µm, which is the sum of two particle sizes in diameter, the composition change was step type, which on the left side of 50 µm was set to the Scheil back diffusion calculation result and on the right side of 50 µm was set to the target composition. The temperature of homogenization heat treatment ranged from 1000 to 1400 • C, maintaining time ranged from 4 to 24 h. The degree of homogenization was represented by the mean value of the standard deviation of each element within the diffusion distance.

Forming Performance
Powder with high sphericity has better fluidity [24], which has an important influence on the powder mixing and distributing uniformly during the SLM process. Therefore, the sphericity and particle size of pure elemental powder were analyzed by backscattering scanning, as shown in Figure 1. The average particle size of four kinds of powders was 37.2 µm, and there are also a few large particles (>53 µm). As shown in Figure 1b, the average particle size of Cr powder was the largest, and that of powders above 10% was higher than 73 µm. The average particle size of Fe powder was the smallest; however, it had a low sphericity. The molten pool shape with different laser parameters was calculated by ANSYS additive science module, and the calculation results are shown in Figure 2. Considering the metallurgical bonding of different elements, the molten pool depth should penetrate at least 1.5 layers thickness (30 µm), and the molten pool width should cover at least the spacing hatch (80 µm). Thermophysical parameters of materials are evitable to get accurate calculation results. Therefore, Thermo-Calc software was used to calculate the average thermal expansion coefficient of the alloy as 1.65 × 10 −5 K −1 . The chamber gas convection coefficient was set to 12.5, and the powder bulk density was set to 0.6 from the database coupled in ANSYS Additive software. The calculation results showed that when the scanning speed varied from 700 to 1100 mm/s and the laser power varied from 100 to 250 W, both the calculated depth and width of molten pools meet the forming requirements. The molten pool shape with different laser parameters was calculated by ANSYS additive science module, and the calculation results are shown in Figure 2. Considering the metallurgical bonding of different elements, the molten pool depth should penetrate at least 1.5 layers thickness (30 μm), and the molten pool width should cover at least the spacing hatch (80 μm). Thermophysical parameters of materials are evitable to get accurate calculation results. Therefore, Thermo-Calc software was used to calculate the average thermal expansion coefficient of the alloy as 1.65 × 10 −5 K −1 . The chamber gas convection coefficient was set to 12.5, and the powder bulk density was set to 0.6 from the database coupled in ANSYS Additive software. The calculation results showed that when the scanning speed varied from 700 to 1100 mm/s and the laser power varied from 100 to 250 W, both the calculated depth and width of molten pools meet the forming requirements. Twenty-four cubes with 10 mm × 10 mm × 7 mm, 9 square bars with 7 mm × 7 mm × 80 mm, 12 sheets with 2 mm × 7 mm × 80 mm, and 1 block with 20 mm × 20 mm × 7 mm were fabricated at one time. The elemental powder and pre-alloying powder were printed separately. All the samples were well formed with good surface quality, as shown in Figure 3. As-printed samples under different printing processes had no visible defects such as pores, cracks, or lack of fusion. The density measurement results of as-printed samples are shown in Figure 4, in which the relative densities are all above 98%. The density of elemental powder samples increased significantly with the increased laser energy from 56 to 104 J/mm 3 , in which the maximum point is 8.06 g/cm 3 (VE = 104 J/mm 3 ). The maximum density of pre-alloyed samples was 8.15 g/cm 3 (VE = 89 J/mm 3 ), and the value fluctuated slightly around 8.1 g/cm 3 with VE varied between 56 and 220 J/mm 3 , which was compatible with the experimental data from Chen Peng et al. [19]. In general, the density of laser in situ alloying FeCoCrNi samples was slightly lower than that of traditional pre-alloyed ones under the same SLM parameters. The density of elemental powder samples would decrease sharply once the energy density is lower than a certain threshold. High-density samples fabricated via elemental powder required a higher laser energy input than the pre-alloyed ones. Twenty-four cubes with 10 mm × 10 mm × 7 mm, 9 square bars with 7 mm × 7 mm × 80 mm, 12 sheets with 2 mm × 7 mm × 80 mm, and 1 block with 20 mm × 20 mm × 7 mm were fabricated at one time. The elemental powder and pre-alloying powder were printed separately. All the samples were well formed with good surface quality, as shown in Figure 3. As-printed samples under different printing processes had no visible defects such as pores, cracks, or lack of fusion. The density measurement results of as-printed samples are shown in Figure 4, in which the relative densities are all above 98%. The density of elemental powder samples increased significantly with the increased laser energy from 56 to 104 J/mm 3 , in which the maximum point is 8.06 g/cm 3 (VE = 104 J/mm 3 ). The maximum density of pre-alloyed samples was 8.15 g/cm 3 (VE = 89 J/mm 3 ), and the value fluctuated slightly around 8.1 g/cm 3 with VE varied between 56 and 220 J/mm 3 , which was compatible with the experimental data from Chen Peng et al. [19]. In general, the density of laser in situ alloying FeCoCrNi samples was slightly lower than that of traditional pre-alloyed ones under the same SLM parameters. The density of elemental powder samples would decrease sharply once the energy density is lower than a certain threshold. High-density samples fabricated via elemental powder required a higher laser energy input than the pre-alloyed ones.

Microstructure Characterization
The metallographic graph (50×) of the XY plane (perpendicular to the printing direc tion) of the as-printed samples fabricated via an elemental powder blend is shown in Fig  ure 5. It can be seen from the figure that the defects such as lack of fusion gradually in creased with the decrease of VE. When the laser energy increased above 104 J/mm 3 , th number of holes basically stayed the same, which was a variation trend that is compatibl with the density. A balling phenomenon occurred in the metallographic structure of sam ples with low energy density. Ruidi Li et al. [16] also observed the same phenomenon from SLM in situ alloying W-10%Cu powders at low energy input density. They also pro posed that the spheroidizing circle was the composition gradient path formed by the lack of diffusing in time. The defeats will deteriorate the mechanical performance of samples so the laser energy for in situ alloying FeCoNiCr alloy should be higher than 104 J/mm 3 .

Microstructure Characterization
The metallographic graph (50×) of the XY plane (perpendicular to the printing direc tion) of the as-printed samples fabricated via an elemental powder blend is shown in Fig  ure 5. It can be seen from the figure that the defects such as lack of fusion gradually in creased with the decrease of VE. When the laser energy increased above 104 J/mm 3 , th number of holes basically stayed the same, which was a variation trend that is compatibl with the density. A balling phenomenon occurred in the metallographic structure of sam ples with low energy density. Ruidi Li et al. [16] also observed the same phenomenon from SLM in situ alloying W-10%Cu powders at low energy input density. They also pro posed that the spheroidizing circle was the composition gradient path formed by the lack of diffusing in time. The defeats will deteriorate the mechanical performance of samples so the laser energy for in situ alloying FeCoNiCr alloy should be higher than 104 J/mm 3 .

Microstructure Characterization
The metallographic graph (50×) of the XY plane (perpendicular to the printing direction) of the as-printed samples fabricated via an elemental powder blend is shown in Figure 5. It can be seen from the figure that the defects such as lack of fusion gradually increased with the decrease of VE. When the laser energy increased above 104 J/mm 3 , the number of holes basically stayed the same, which was a variation trend that is compatible with the density. A balling phenomenon occurred in the metallographic structure of samples with low energy density. Ruidi Li et al. [16] also observed the same phenomenon from SLM in situ alloying W-10%Cu powders at low energy input density. They also proposed that the spheroidizing circle was the composition gradient path formed by the lack of diffusing in time. The defeats will deteriorate the mechanical performance of samples, so the laser energy for in situ alloying FeCoNiCr alloy should be higher than 104 J/mm 3 . The microstructure and composition mapping results on the XZ plane (parallel t printing direction) of elemental powder samples with a magnification of 350× are sh in Figure 6. The molten pool shape of the samples with the maximum laser energy i (223 J/mm 3 ) were equal in size, and the boundaries were clearly visible, as shown in Fi 6a. Obvious spheroidal areas appeared in the scanning region when VE was reduce 173 J/mm 3 (Figure 6b), and the depth of the molten pool also decreased. The resul energy spectrum scanning of those spheroidal areas are shown in Figure 6e. The ran distribution of the spherical areas featured pure Cr-rich zones. When VE (Figure went down further, the depth of the molten pool decreased synchronously, the mo pool boundary gradually became blurred, and the Cr-rich zones that were mainly dis uted in the laps between molten pools increased greatly. The boundaries between th rich zones and matrix were well fused, and there was no hole or lack of fusion. Fu element composition scanning and microstructure observation of HEE1 samples shown in Figure 7a. Except for a few occasional Cr-rich zones, other elements were ho geneously distributed. Four micro-areas of about 0.3 mm × 0.3 mm were randoml lected on the XZ plane for composition line scanning. The average value of the diag scanning data of micro-areas is used to characterize the average element composition results are shown in Table 3. Compared with the target components, the average devia value of Fe and Co elements from the target components is only 2%, while the devia value of Cr and Ni elements is less than 13%. The micro-area was scanned parallel t printing direction with a length of about 260 μm, as shown in Figure 7b. Significant position fluctuations between layers can be observed. The HEE5 sample (VE = 173 J/m was characterized in the same way as the results shown in Table 3. Due to the decrea laser energy density, growing pure Cr-rich regions appeared in the sample, and the m imum element composition deviated to 17%. Semi-quantitative composition analys HEE7 (VE = 104 J/mm 3 ) by spark spectroscopy showed that the deviation of the hig The microstructure and composition mapping results on the XZ plane (parallel to the printing direction) of elemental powder samples with a magnification of 350× are shown in Figure 6. The molten pool shape of the samples with the maximum laser energy input (223 J/mm 3 ) were equal in size, and the boundaries were clearly visible, as shown in Figure 6a. Obvious spheroidal areas appeared in the scanning region when VE was reduced to 173 J/mm 3 (Figure 6b), and the depth of the molten pool also decreased. The results of energy spectrum scanning of those spheroidal areas are shown in Figure 6e. The random distribution of the spherical areas featured pure Cr-rich zones. When VE (Figure 6c,d) went down further, the depth of the molten pool decreased synchronously, the molten pool boundary gradually became blurred, and the Cr-rich zones that were mainly distributed in the laps between molten pools increased greatly. The boundaries between the Cr-rich zones and matrix were well fused, and there was no hole or lack of fusion. Further element composition scanning and microstructure observation of HEE1 samples are shown in Figure 7a. Except for a few occasional Cr-rich zones, other elements were homogeneously distributed. Four micro-areas of about 0.3 mm × 0.3 mm were randomly selected on the XZ plane for composition line scanning. The average value of the diagonal scanning data of micro-areas is used to characterize the average element composition. The results are shown in Table 3. Compared with the target components, the average deviation value of Fe and Co elements from the target components is only 2%, while the deviation value of Cr and Ni elements is less than 13%. The micro-area was scanned parallel to the printing direction with a length of about 260 µm, as shown in Figure 7b. Significant composition fluctuations between layers can be observed. The HEE5 sample (VE = 173 J/mm 3 ) was characterized in the same way as the results shown in Table 3. Due to the decrease of laser energy density, growing pure Cr-rich regions appeared in the sample, and the maximum element composition deviated to 17%. Semi-quantitative composition analysis of HEE7 (VE = 104 J/mm 3 ) by spark spectroscopy showed that the deviation of the highest element composition in the radius of 3 mm was less than 20%.      The phase diagram of FeCoCrNi equiatomic HEAs was calculated by Thermo-Calc software, as shown in Figure 8. The melting point of the homogeneous solid was about 1450 • C, while the single-phase FCC solid solution can be formed in the temperature range of 700-1400 • C. When the components of Fe, Co, and Cr deviated in the range of 0-20%, the matrix phase structure remained FCC unchanged, and the sigma phase precipitated in the matrix phase as the temperature went down. The matrix phase changed from a single-phase FCC structure to a two-phase FCC + BCC structure when the mass percentage of element Ni was less than 3%. In general, the deviation of element composition is all within 20% according to the above composition characterization results. The matrix phase structure of the as-printed samples should be an FCC structure according to the calculated phase diagram analysis. The XRD structures of the laser in situ alloying samples and pre-alloyed powder samples were analyzed. As shown in Figure 9, the structure of the laser in situ alloying sample and the pre-alloyed sample is an FCC single phase structure, and the phase structure is consistent with that of the sample prepared by the traditional method, which proves that the four element powders have completed the micro alloying during SLM printing. The characteristics of the SLM process determine that the sample will have serious solidification segregation. When the cooling rate of the SLM printing process is 10 5 • C/s, the Scheil back diffusion calculation results of the composition segregation between the first solidified dendrite arm and the later solidified dendrite arm in the microstructure are as shown in Figure 10. FCC_L12 is the matrix phase with a crystalline structure of ordered face-centered cubic. Affected by the interdiffusion between elements, the segregation of the Cr element is the most serious (32%), followed by the Co element (16%), while the rapid solidification process has little effect on the microstructure segregation of Fe and Ni elements. The calculation results are consistent with the measured composition deviation law of the laser in situ alloying sample of elemental powder, indicating that one of the factors leading to the composition deviation is caused by the extreme unbalanced solidification process. 21, 11, x FOR PEER REVIEW 11 of 17         Although the micro-segregation brought by the SLM process can be impro increasing the laser power and spot diameter alongside reducing the scanning spe inevitable that the subsequent high-temperature heat treatment must be combine prove the micro-segregation. Therefore, the homogenization diffusion model built TRA was used to calculate the effect of the post-heat treatment regime on element enization. The calculation result is shown in Figure 11. The homogenization of seg elements could be achieved by using the heat treatment temperature of 1400 °C for time. The homogenization of segregated elements can be largely achieved by us heat treatment temperature of 1200 °C for more than 12 h. There was little effec elements' diffusion even if the as-printed samples were under the heat treatment ature of 1000 °C for more than 24 h. Long time and high temperature heat treatme lead to grain boundary melting (over burning) or grain coarsening (over heating); treatment schedule of 1200 °C for 20 h was selected for the experiment. The microst energy spectrum scanning results of the samples after heat treatment are shown in 12. The elements were evenly distributed, and the grain morphology was a mixin ture of columnar grain and equiaxed grain. Therefore, the homogenization heat tre schedule of laser in situ alloying high-entropy alloy samples can be accurately o by using the Scheil model coupled with the DICTRA model. Although the micro-segregation brought by the SLM process can be improved by increasing the laser power and spot diameter alongside reducing the scanning speed, it is inevitable that the subsequent high-temperature heat treatment must be combined to improve the micro-segregation. Therefore, the homogenization diffusion model built in DICTRA was used to calculate the effect of the post-heat treatment regime on element homogenization. The calculation result is shown in Figure 11. The homogenization of segregated elements could be achieved by using the heat treatment temperature of 1400 • C for a short time. The homogenization of segregated elements can be largely achieved by using the heat treatment temperature of 1200 • C for more than 12 h. There was little effect on the elements' diffusion even if the as-printed samples were under the heat treatment temperature of 1000 • C for more than 24 h. Long time and high temperature heat treatment may lead to grain boundary melting (over burning) or grain coarsening (over heating); the heat treatment schedule of 1200 • C for 20 h was selected for the experiment. The microstructure energy spectrum scanning results of the samples after heat treatment are shown in Figure 12. The elements were evenly distributed, and the grain morphology was a mixing structure of columnar grain and equiaxed grain. Therefore, the homogenization heat treatment schedule of laser in situ alloying high-entropy alloy samples can be accurately obtained by using the Scheil model coupled with the DICTRA model.

Mechanical Properties
The tensile strength of as-printed samples is shown in Figure 13. The strength of elemental powder samples was equal to that of pre-alloyed powder samples under the optimal energy input (140 J/mm 3 ). When VE varied between 69 and 173 J/mm 3 , the tensile strength of elemental powder samples increased along with the energy density, and the value ranged from 570 to 680 MPa. The tensile strength of pre-alloy powder samples decreased with the increase of energy density, and the value ranged from 610 to 680 MPa. The average hardness of the elemental powder samples was slightly lower than that of the pre-alloyed powder samples, as shown in Figure 14. The hardness of the elemental

Mechanical Properties
The tensile strength of as-printed samples is shown in Figure 13. The strength of ele mental powder samples was equal to that of pre-alloyed powder samples under the opt mal energy input (140 J/mm 3 ). When VE varied between 69 and 173 J/mm 3 , the tensil strength of elemental powder samples increased along with the energy density, and th value ranged from 570 to 680 MPa. The tensile strength of pre-alloy powder samples de creased with the increase of energy density, and the value ranged from 610 to 680 MPa The average hardness of the elemental powder samples was slightly lower than that o the pre-alloyed powder samples, as shown in Figure 14. The hardness of the elementa

Mechanical Properties
The tensile strength of as-printed samples is shown in Figure 13. The strength of elemental powder samples was equal to that of pre-alloyed powder samples under the optimal energy input (140 J/mm 3 ). When VE varied between 69 and 173 J/mm 3 , the tensile strength of elemental powder samples increased along with the energy density, and the value ranged from 570 to 680 MPa. The tensile strength of pre-alloy powder samples decreased with the increase of energy density, and the value ranged from 610 to 680 MPa.
The average hardness of the elemental powder samples was slightly lower than that of the pre-alloyed powder samples, as shown in Figure 14. The hardness of the elemental powder samples has no obvious correlation with energy density when it ranged from 210 to 240 HV. The hardness of the pre-alloyed powder samples increased slightly along with the energy density, which ranged from 220 to 260 HV. The comparison of mechanical properties of FeCoCrNi alloy prepared by different methods as reported in other studies in the literature are listed in Table 4. Brif et al. [25] used laser selective melting technology to print FeCoCrNi HEAs for the first time in 2015 and confirmed the printability of the alloy system. In recent years, the alloy system has also been fabricated by extrusion, rolling, and other hot working processes and obtained excellent properties [19,[26][27][28]. According to the literature data, the strength and hardness of the elemental powder samples under an optimized process were comparable to those prepared by the pre-alloying powder and much higher than those prepared by the traditional arc melting method.
Metals 2021, 11, x FOR PEER REVIEW powder samples has no obvious correlation with energy density when it ranged f to 240 HV. The hardness of the pre-alloyed powder samples increased slightly alo the energy density, which ranged from 220 to 260 HV. The comparison of me properties of FeCoCrNi alloy prepared by different methods as reported in othe in the literature are listed in Table 4. Brif et al. [25] used laser selective melting tec to print FeCoCrNi HEAs for the first time in 2015 and confirmed the printabili alloy system. In recent years, the alloy system has also been fabricated by extrus ing, and other hot working processes and obtained excellent properties [19,26− cording to the literature data, the strength and hardness of the elemental powder under an optimized process were comparable to those prepared by the pre-alloyi der and much higher than those prepared by the traditional arc melting method.  powder samples has no obvious correlation with energy density when it ran to 240 HV. The hardness of the pre-alloyed powder samples increased sligh the energy density, which ranged from 220 to 260 HV. The comparison properties of FeCoCrNi alloy prepared by different methods as reported in in the literature are listed in Table 4. Brif et al. [25] used laser selective melti to print FeCoCrNi HEAs for the first time in 2015 and confirmed the prin alloy system. In recent years, the alloy system has also been fabricated by e ing, and other hot working processes and obtained excellent properties [ cording to the literature data, the strength and hardness of the elemental po under an optimized process were comparable to those prepared by the preder and much higher than those prepared by the traditional arc melting me

Discussion
There were three main factors affecting the composition deviation of elemental powder samples: (1) composition deviations caused by uneven powder mixing, powder fluidity, and powder size difference during the powder laying process of SLM; (2) element segregation caused by the rapid solidification process; and (3) incomplete alloying of high melting point elements caused by insufficient laser energy density. The inhomogeneity of the powder blend in the paving process will directly lead to a great composition deviation of the final specimen from the target component. In this experiment, batch mixing powder was used, which resulted in up to 20% regional component deviation. Regional component deviation can be significantly improved if micro-zone real-time powder mixing was adopted. Serious element microstructure segregation caused by rapid melting and solidification can be eliminated by the subsequent high-temperature heat treatment.
The main reasons for the unalloyed chromium enrichment zone in the as-printed samples were as follows. (1) The melting point of Cr powder is the highest: at least 340 • C higher than other elements. (2) The average particle size of Cr powder is larger than that of other powders; more than 10% of the particle diameter is greater than 73 µm, which exceeds the recommended particle size of SLM. (3) The unalloyed zone mainly occurred in the overlap position of the molten pool, that is, the laser energy density in the heat-affected zone is not enough to realize the alloying of elements. As the scanning strategy of SLM was 80 microns hatch spacing and 80 µm laser spot diameter; that is, the overlap rate is 0, so there is a heat source blind zone at the bottom of proximate molten pools. Since the melting point of Cr is the highest, FeCoNi acted as a solution, Cr acted as a solute, and laser in situ alloying was a liquid phase sintering process. The melted FeCoNi liquid phase continuously moistened the Cr element, so the interface between the Cr-rich region and matrix was well fused, and no holes were formed. However, due to the insufficient laser energy, rapid diffusion time (only ms level), and the large particle size of the Cr powder will lead to an incomplete Cr diffusion in the liquid phase. In addition, element diffusion in the solid phase was very slow and negligible, so the Cr-rich region remained heterogeneous to room temperature after multiple thermal cycles. Overall, the alloy elements of laser in situ alloying technology should be selected as close as possible to the melting point. It is necessary to improve the laser energy density and increase the overlap rate of the molten pool in the printing process design if there is a high melting point element.

Conclusions
Bulk samples of FeCoCrNi equiatomic HEAs with good one-time forming performance were prepared by the SLM in situ alloying method with the processing optimization via finite element analysis and thermodynamic calculations. Combined with a real-time quantitative powder feeding device, the technology scheme can be extended to the highthroughput preparation of novel HEAs to obtain large amounts of data (compositionprocess-performance) for material design efficiently and quickly. The microstructure and properties of SLM in situ alloying FeCoCrNi samples via elemental powders are as follows.
(1) The SLM in situ alloying of pure elements (Fe, Co, Cr, Ni) was achieved completely within the reasonable SLM process parameter (VE = 224 J/mm 3 ) according to the experimental analysis and calculation results. The deviation of as-printed samples from the target composition can be controlled within 20 wt %. The matrix phase was an ordered FCC structure, which was the same with pre-alloyed powder samples. Insufficient energy, the low overlap rate of the molten pool, and the large particle size of powders led to incomplete alloying of the high melting point powder.
(2) There was a linear relationship between the density and laser energy input of elemental powder samples. It is only 1% lower than that of pre-alloyed powder samples when the energy density exceeds the critical value (VE = 224 J/mm 3 ). The strength and hardness of the samples were comparable to those of the pre-alloyed powder samples, which are much higher than those prepared by the traditional arc melting method. To obtain excellent performance, a higher energy input of element powder samples was required than that of pre-alloyed powder.
(3) The microstructure segregation of as-printed samples was eliminated by hightemperature heat treatment. The SHEIL + DICTRA calculation model can quantitatively give an appropriate heat treatment schedule. The calculation results show that the complete homogenization of elemental powder samples can be achieved by using the heat treatment schedule of 1400 • C for 4 h or 1200 • C for 20 h. The experimental results show that homogeneous samples can be obtained by using the heat treatment temperature of 1200 • C for 20 h.

Data Availability Statement:
The data presented in this study are available on request from the corresponding author. The data are not publicly available due to potential commercial sensitivity.