Control of Crystallographic Texture and Mechanical Properties of Hastelloy-X via Laser Powder Bed Fusion

: The inﬂuence of various laser powder bed fusion (LPBF) process parameters on the crystallographic textures and mechanical properties of a typical Ni-based solid-solution strengthened alloy, Hastelloy-X, was examined. Samples were classiﬁed into four groups based on the type of crystallographic texture: single crystalline-like microstructure with <100>//build direction (BD) (<100>-SCM), single crystalline-like microstructure with <110>//BD (<110>-SCM), crystallographic lamellar microstructure (CLM), or polycrystalline microstructure (PCM). These four crystallographic textures were realized in Hastelloy-X for the ﬁrst time here to the best of our knowledge. The mechanical properties of the samples varied depending on their texture. The tensile properties were affected not only by the Schmid factor but also by the grain size and the presence of lamellar boundaries (grain boundaries). The lamellar boundaries at the interface between the <110>//BD oriented main layers and the <100>//BD-oriented sub-layers of CLM contributed to the resistance to slip transmission and the increased proof stress. It was possible to control a wide range of crystallographic microstructures via the LPBF process parameters, which determines the melt pool morphology and solidiﬁcation behavior.

In terms of Ni-based superalloys, the LPBF process has been investigated for various alloys with different strengthening mechanisms, such as the solid-solution strengthened alloy Hastelloy-X [28], the γ" (typically Ni 3 Nb) precipitation strengthened alloy IN718 [15], and the γ' (typically Ni 3 Al) precipitation strengthened alloy IN939 [29]. The major applications of Ni-based superalloys include aircraft engines and industrial gas turbines. Generally, solid-solution strengthened alloys are widely used in combustor components that require good formability for complex shapes and oxidation resistance [30]. γ" precipitation strengthened alloys are utilized for many components, such as shafts and disks, which are not directly exposed to combustion gases and require excellent mechanical properties in the medium-temperature range, up to 700 • C [31]. γ precipitation strengthened alloys are generally utilized for turbine components that require excellent creep properties [32].
Solid-solution strengthened alloys are useful for practical applications for the following reasons: (1) the required mechanical properties are less than that of precipitation strengthened alloys, (2) they do not require various steps of heat treatment, and (3) the LPBF process can manufacture complex parts which are useful for their application in combustor components. For practical use, it is necessary to investigate a wide range of process parameters and determine the appropriate process window. However, previous studies have focused only on examining the effects of a limited number of process parameters on the microstructure and mechanical properties [33][34][35], understanding precipitates and recrystallization during heat treatment [36,37], analyzing the factors that cause micro-cracks [38][39][40][41], and improving mechanical properties by dispersion of TiC particles [42]. Therefore, a sufficient understanding of the effects of a wider range of process parameters on crystallographic textures and their relation to mechanical properties has not been established.
In this study, we examined the influence of a wide range of process parameters on the crystallographic textures and mechanical properties of a typical solid-solution strengthened alloy, Hastelloy-X. The microstructural evolution is also discussed by focusing on the melt pool shape, which is a microscopic melting unit, in order to obtain a guideline for determining the optimal process parameters for obtaining desired microstructural characteristics.

Fabrication of Hastelloy-X Samples via LPBF Process
Gas-atomized Hastelloy-X powder (EOS, Krailling, Germany) with the chemical composition included in Table 1 was used in this study. A field-emission scanning electron microscopy (FE-SEM; JIB-4610F, JEOL, Tokyo, Japan) image of the Hastelloy-X powder is displayed in Figure 1a. The particle size distribution was determined as D 10 = 17.4 µm, D 50 = 33.9 µm, and D 90 = 63.1 µm, using a Mastersizer 3000 (Malvern Panalytical, Malvern, UK). The flowability, measured by a Revolution (Mercury Scientific, Newtown, CT, USA), was 14.68 mJ/kg for the avalanche energy, 45.1 • for the avalanche angle, 34.3 • for the rest angle, and 1.87 for the surface fractal. An EOS M 290 (EOS, Krailling, Germany) was used for LPBF, and samples with dimensions of 10 mm (x direction), 10 mm (y direction), and 50 mm (z direction; building direction (BD)) ( Figure 1b) were produced. The building platform was heated to 80 • C, and the building atmosphere was maintained at an oxygen content of less than 100 ppm by Ar gas. The layer thickness t L was 0.06 mm, the laser power P was 180-360 W (in 60 W steps), the scan speed v was 600-1400 mm/s (in 200 mm/s step), and the hatch distances d were 0.08 and 0.10 mm. The volumetric energy density E vol [J/mm 3 ] can be calculated using Equation (1). Two scan strategies, X scan and XY scan [22], were employed. To investigate a wide range of process parameters, 80 parameter conditions were considered.
composition included in Table 1 was used in this study. A field-e tron microscopy (FE-SEM; JIB-4610F, JEOL, Tokyo, Japan) image o der is displayed in Figure 1a. The particle size distribution was de μm, D50 = 33.9 μm, and D90 = 63.1 μm, using a Mastersizer 3000 (Ma vern, UK). The flowability, measured by a Revolution (Mercury Sc USA), was 14.68 mJ/kg for the avalanche energy, 45.1° for the ava the rest angle, and 1.87 for the surface fractal.

Microstructure Characterization
The y-z cross-sections were cut out, embedded in resins, and polished with emery paper up to #4000 and colloidal silica carried out as part of the initial preparation for microstructure observation. An optical microscope (OM; BX60, Olympus, Tokyo, Japan) was used to observe internal defects. FE-SEM observations were conducted to determine the shape of the melt pool after etching with a HF:HNO 3 :H 2 O = 35:55:100 solution at 50-60 • C for 90 min. Electron backscatter diffraction (EBSD), which the FE-SEM was equipped for, was used to identify the crystallographic textures under an acceleration voltage of 20 kV, a current of 16 nA, and a step size of 2 µm at 100× magnification. Inverse pole figure (IPF) maps, pole figures, grain boundary maps, and Euler angles were obtained using analysis software (AztecHKL 3.0, Oxford Instruments, High Wycombe, UK). For quantitative evaluation, the degree of orientation p hkl with respect to the observation direction <hkl> was obtained using the Euler angles of every measurement point analyzed by the HKL Channel 5 software (Oxford Instruments, High Wycombe, UK), according to Equation (2). Here, α hkl is the smallest angle between the observation direction and the equivalent orientations of <hkl>.

Mechanical Property Tests
The Young's modulus was determined using the free resonance elastic modulus measuring method (JE2-C1/JE2-RT, Nihon Techno-Plus, Osaka, Japan), using small samples cut to 2 mm thickness (x direction), 8 mm width (y direction), and 40 mm length (z direction) from the built samples. The Young's modulus E [GPa] was calculated using Equation (3), employing the measured resonance frequency f [Hz], the specimen length l [mm], specimen thickness t s [mm], and specimen density ρ [g/cm 3 ]. The density was measured using the Archimedes method with an analytical balance (AUX 220, Shimadzu, Kyoto, Japan).
Tensile samples were cut to 0.8 mm thickness (x direction), 2 mm width (y direction), and 5 mm gauge length (z direction) from the built samples, with tensile properties obtained using an autograph (AGX-V, Shimadzu, Kyoto, Japan) and analyzed using TrapeziumX-V software (Shimadzu, Kyoto, Japan). The initial strain rate was set to 1%/min with the test environment at room temperature in a vacuum atmosphere.

Microstructure Characterization
The effect of E vol on the density of the samples is displayed in Figure 2 including representative OM images of the y-z cross-section. The highest density was observed at E vol = 50 J/mm 3 . When E vol was less than 35 J/mm 3 , the density decreased rapidly and lack-of-fusion defects were observed. When E vol exceeded approximately 95 J/mm 3 , the density decreased more gradually, and keyhole-type spherical defects (gas porosities) were observed. These two types of defects are widely recognized as part of the LPBF process [43]. Based on these preliminary experiments, the optimum process parameter determined for densification of the samples of Hastelloy-X was E vol in the range of 35-95 J/mm 3 .
Tensile samples were cut to 0.8 mm thickness (x direction), 2 mm width (y direction), and 5 mm gauge length (z direction) from the built samples, with tensile properties obtained using an autograph (AGX-V, Shimadzu, Kyoto, Japan) and analyzed using Trape-ziumX-V software (Shimadzu, Kyoto, Japan). The initial strain rate was set to 1%/min with the test environment at room temperature in a vacuum atmosphere.

Microstructure Characterization
The effect of Evol on the density of the samples is displayed in Figure 2 including representative OM images of the y-z cross-section. The highest density was observed at Evol = 50 J/mm 3 . When Evol was less than 35 J/mm 3 , the density decreased rapidly and lackof-fusion defects were observed. When Evol exceeded approximately 95 J/mm 3 , the density decreased more gradually, and keyhole-type spherical defects (gas porosities) were observed. These two types of defects are widely recognized as part of the LPBF process [43]. Based on these preliminary experiments, the optimum process parameter determined for densification of the samples of Hastelloy-X was Evol in the range of 35-95 J/mm 3 . IPF maps and pole figures under typical process parameters with d = 0.10 mm for the X-scan strategy are displayed in Figure 3 with the results for Evol < 35 J/mm 3 and Evol > 95 J/mm 3 indicated with crosses. The grains elongated along the BD were observed for a wide range of process parameters. Additionally, the crystallographic textures strongly depended on the process parameters. Subsequently, as illustrated in Figure   IPF maps and pole figures under typical process parameters with d = 0.10 mm for the X-scan strategy are displayed in Figure 3 with the results for E vol < 35 J/mm 3 and E vol > 95 J/mm 3 indicated with crosses. The grains elongated along the BD were observed for a wide range of process parameters. Additionally, the crystallographic textures strongly depended on the process parameters. Subsequently, as illustrated in Figure 3A

Effect of Crystallographic Texture on the Young's Modulus
The Young's modulus was measured to investigate the effect of the crystallographic textures obtained by the LPBF process on the sample's mechanical properties. The Young's modulus of a single crystal with respect to an arbitrary orientation < > is obtained using Equation (4)

Effect of Crystallographic Texture on the Young's Modulus
The Young's modulus was measured to investigate the effect of the crystallographic textures obtained by the LPBF process on the sample's mechanical properties. The Young's modulus of a single crystal with respect to an arbitrary orientation <lmn> is obtained using Equation (4) with the elastic stiffness constants c 11 , c 12 , and c 44 . As c 11 , c 12 , and c 44 of Hastelloy-X are 230.40, 156.12, and 121.77 GPa, respectively [44], the theoretical Young's modulus of the <100>, <110> and <111> orientations were calculated as 104.28, 203.62, and 298.36 GPa, respectively.
The two methods of averaging the Young's modulus of polycrystallines are known as the Voigt average E V (Equation (5)) [45], assuming the same strain state of each crystal, Crystals 2021, 11, 1064 6 of 12 and the Reuss average E R (Equation (6)) [46], assuming the same stress state of each crystal. Then, the Voigt-Reuss-Hill approximation E H (Equation (7)) [47], taking the average of E V and E R , is typically employed. Here, s 11 , s 12 , and s 44 are the elastic compliance constants, calculated using c 11 , c 12 , and c 44 . Therefore, the theoretical Young's modulus of polycrystalline Hastelloy X was calculated as 199.01 GPa.
The relationship between the degree of orientation and the measured Young's modulus is illustrated in Figure 4. The maximum value of p hkl is 1 for a single crystal whose orientation coincides perfectly with the observation direction <hkl>. In a completely random polycrystalline structure, p 100 is 0.701, and p 110 is 0.835 [48]. It was confirmed that the Young's modulus decreases as p 100 increases, and that <100>-SCM has a theoretical Young's modulus in the <100> orientation of 104.28 GPa. Similarly, the Young's modulus increased as p 110 increased, and <110>-SCM demonstrated a Young's modulus of 203.62 GPa, close to the theoretical value. Therefore, the crystallographic textures uniquely obtained via the LPBF process demonstrated a Young's modulus comparable to the theoretical value, corresponding to the orientation. That is, desired mechanical properties may be obtained by controlling the process parameters.
The two methods of averaging the Young's modulus of polycrystallines are known as the Voigt average EV (Equation (5)) [45], assuming the same strain state of each crystal, and the Reuss average ER (Equation (6)) [46], assuming the same stress state of each crystal. Then, the Voigt-Reuss-Hill approximation EH (Equation (7)) [47], taking the average of EV and ER, is typically employed. Here, , , and are the elastic compliance constants, calculated using , , and . Therefore, the theoretical Young's modulus of polycrystalline Hastelloy X was calculated as 199.01 GPa.
The relationship between the degree of orientation and the measured Young's modulus is illustrated in Figure 4. The maximum value of 〈 〉 is 1 for a single crystal whose orientation coincides perfectly with the observation direction <hkl>. In a completely random polycrystalline structure, 〈 〉 is 0.701, and 〈 〉 is 0.835 [48]. It was confirmed that the Young's modulus decreases as 〈 〉 increases, and that <100>-SCM has a theoretical Young's modulus in the <100> orientation of 104.28 GPa. Similarly, the Young's modulus increased as 〈 〉 increased, and <110>-SCM demonstrated a Young's modulus of 203.62 GPa, close to the theoretical value. Therefore, the crystallographic textures uniquely obtained via the LPBF process demonstrated a Young's modulus comparable to the theoretical value, corresponding to the orientation. That is, desired mechanical properties may be obtained by controlling the process parameters.

Impact of Crystallographic Textures on Tensile Properties
The true stress-true strain curves of the four groups of crystallographic textures are shown in Figure 5. Additionally, Table 2 summarizes the mean values of the tensile properties with standard deviations.
The Schmid factors (SF) for the loading axes <100> and <110> were both 0.408. As CLM can be regarded as a composite microstructure consisting of main layers with <110>//BD orientation and sub-layers with <100>//BD orientation, the theoretical SF is also 0.408. Here, the 0.2% proof stresses of <100>-SCM and <110>-SCM were 419.2 ± 3.5 and 419.9 ± 13.5 MPa, respectively, which is consistent with the identical SFs. The higher work-hardening rate for <100>-SCM may be due to the higher frequency of dislocation pileup associated with multiple slips, as the loading axis parallel to <100> is not accompanied by crystal rotation, but by an 8-fold multiple slip during tensile testing [49,50]. However, Crystals 2021, 11, 1064 7 of 12 the 0.2% proof stress of CLM was 434.1 ± 12.2 MPa, higher than that of the <100>-SCM and <110>-SCM, despite the theoretically identical SF. PCM demonstrated the highest 0.2% proof stress of 493.1 ± 13.3 MPa. As illustrated in the grain boundary maps of Figure 6d, it was determined that the proof stress obeys the Hall-Petch relationship.
The required values of 0.2% proof stress, UTS, and elongation for cast Hast specified in AMS5390 are 240 MPa, 379 MPa, and 8%, respectively ( Table 2). Thos Hastelloy-X sheets specified in ASTM B435-06 are 240 MPa, 655 MPa, and 35%, tively ( Table 2). The tensile properties of all four groups exceeded the specificat AMS5390, indicating the feasibility of part manufacturing utilizing the crystallog textures unique to the LPBF process. With respect to ASTM B435-06, PCM satisf requirements for all specified values. By changing the crystallographic orientati grain boundary density via fabrication using LPBF, the proof stress, UTS, and elon can be chosen according to the mechanical requirements.      The required values of 0.2% proof stress, UTS, and elongation for cast Hastelloy-X specified in AMS5390 are 240 MPa, 379 MPa, and 8%, respectively (Table 2). Those  Table 2). The tensile properties of all four groups exceeded the specifications of AMS5390, indicating the feasibility of part manufacturing utilizing the crystallographic textures unique to the LPBF process. With respect to ASTM B435-06, PCM satisfied the requirements for all specified values. By changing the crystallographic orientation and grain boundary density via fabrication using LPBF, the proof stress, UTS, and elongation can be chosen according to the mechanical requirements.

Microstructure Evolution during LPBF
Among the four groups of crystallographic textures obtained in this study, <100>-SCM, CLM, and <110>-SCM, which were presumed to have been obtained by characteristic solidification phenomena in the LPBF process, are discussed with a focus on the melt pool, which is a basic unit of solidification.
The SEM images of the melt pool visualized by etching are displayed in Figure 7. Fine cells were observed in the melt pool, indicating that cell growth occurred in all cases. The elongation direction of the cells is known to coincide with <100>, which is the prior growth direction of FCC metals. Here, a remarkable difference was observed in the elongation direction of the cells. Those in the <100>-SCM developed from the bottom of the melt pool in the BD and from the side of the melt pool in the horizontal direction. However, in <110>-SCM, cell growth on a 45 • incline to the BD was observed from the bottom to the side of the melt pool. Regarding CLM, the main layers in which the cells grow with a 45 • tilt to the BD from the side of the melt pool, and the sublayers in which the cells grow parallel to the BD from the bottom of the melt pool, are alternately arranged.

Expectations for the Crystallographic Textures on Mechanical and Oxidation Properties
Expectations for the crystallographic texture control in Hastelloy-X are described in terms of the mechanical properties and oxidation resistance.
The tensile properties of the PCM obtained in this study were approximately equivalent to those reported in the literature [32,35,36,39]. However, this study reports on the <100>-SCM, CLM, and <110>-SCM of Hastelloy-X developed via the LPBF process for the first time. Among these, the <100> orientation in single-crystalline Ni-based superalloys is important not only for its excellent creep properties as compared with other orientations [29], but also due to its excellent low-cycle fatigue properties [51]. Furthermore, it has been reported that in pure chromium (Cr) developed via the LPBF process, <100>-oriented Cr demonstrated superior oxidation resistance compared to randomly oriented Cr [52]. This is explained by the anisotropy of the diffusion coefficient of Cr2O3. Hastelloy-X, which forms the same Cr2O3 oxide film, could exhibit a similar advantage based on the crystallographic texture.
The 0.2% proof stress of CLM was significantly higher than that of SCM, despite the identical SF, with the presence of a lamellar boundary being the microstructural difference between the CLM and SCM. This lamellar boundary consists of grain boundaries with an angular difference of 45 ° and is likely to be the source of the higher proof stress. As proposed for bicrystals, the boundary between single crystals can be a resistant to slip transmission [53]. The ideal transmission coefficient calculated at the CLM boundary was 0.82 [26], where the value for a single crystal is 1. It was suggested that the transmission coefficient is less than 1, but moderately high, which gives the CLM good elongation and The relationship between the cell growth direction and melt pool shape was examined in detail. As a representative value for the shapes of the melt pool, the radius of curvature r [µm] at the melt pool bottom was calculated according to Equation (8) [25]. Here, the shapes of the melt pool in the y-z plane were approximated by a function f (a). The calculated radii of curvature were 61.1 ± 9.41, 50.9 ± 6.56, and 39.2 ± 2.96 µm for <100>-SCM, CLM, and <110>-SCM, respectively. As the bottom of the melt pool became flatter (r increased), the <100>//BD growth became dominant, resulting in <100>-SCM evolution.

Expectations for the Crystallographic Textures on Mechanical and Oxidation Properties
Expectations for the crystallographic texture control in Hastelloy-X are described in terms of the mechanical properties and oxidation resistance.
The tensile properties of the PCM obtained in this study were approximately equivalent to those reported in the literature [32,35,36,39]. However, this study reports on the <100>-SCM, CLM, and <110>-SCM of Hastelloy-X developed via the LPBF process for the first time. Among these, the <100> orientation in single-crystalline Ni-based superalloys is important not only for its excellent creep properties as compared with other orientations [29], but also due to its excellent low-cycle fatigue properties [51]. Furthermore, it has been reported that in pure chromium (Cr) developed via the LPBF process, <100>-oriented Cr demonstrated superior oxidation resistance compared to randomly oriented Cr [52]. This is explained by the anisotropy of the diffusion coefficient of Cr 2 O 3 . Hastelloy-X, which forms the same Cr 2 O 3 oxide film, could exhibit a similar advantage based on the crystallographic texture.
The 0.2% proof stress of CLM was significantly higher than that of SCM, despite the identical SF, with the presence of a lamellar boundary being the microstructural difference between the CLM and SCM. This lamellar boundary consists of grain boundaries with an angular difference of 45 • and is likely to be the source of the higher proof stress. As proposed for bicrystals, the boundary between single crystals can be a resistant to slip transmission [53]. The ideal transmission coefficient calculated at the CLM boundary was 0.82 [26], where the value for a single crystal is 1. It was suggested that the transmission coefficient is less than 1, but moderately high, which gives the CLM good elongation and higher proof stress compared to SCM.
As Hastelloy-X is a solid-solution strengthened Ni-based superalloy, solution treatments and aging treatments are not necessarily required. That is, Hastelloy-X can be used without recrystallization, while also utilizing crystallographic textures uniquely created by the LPBF process. Therefore, improving our understanding of the textures of Hastelloy-X was important, and the utilization of the LPBF process should be more widely adopted in the future based on the relationship between process parameters, crystallographic textures, and mechanical properties.

Conclusions
In this study, we examined the influence of a wide range of process parameters on the crystallographic textures and mechanical properties of a typical solid-solution strengthened alloy, Hastelloy-X. The following conclusions were drawn: (1) Among a wide range of process parameters, dense samples were obtained at a volumetric energy density of 35-95 J/mm 3 . They were classified into four groups, <100>-SCM, CLM, <110>-SCM, and PCM, based on the crystallographic textures. (2) The Young's modulus of the samples constructed via the LPBF process showed anisotropy close to the theoretical value depending on the degree of orientation. It was suggested that the desired mechanical properties could be obtained by controlling the process parameters. (3) The tensile properties were affected not only by the SF, but also by the grain size and the presence of grain boundaries (lamellar boundaries). PCM with smaller grains showed a higher 0.2% proof stress than the other groups. Grain boundaries at the interface between the <110>//BD main layers and the <100>//BD sub-layers of CLM played a role in the resistance to slip transmission and contributed to the increase in proof stress. (4) By controlling the radius of curvature at the bottom of the melt pool using the process parameters, it is possible to control a wide range of microstructures, including <100>-SCM, <110>-SCM, and CLM, which is a microstructure uniquely obtained via the LPBF process.
It was demonstrated that it is possible to control the crystallographic textures using the LPBF process by appropriately choosing the process parameters, and the mechanical properties corresponding to the crystallographic textures can be designed and obtained. Microstructure control is expected to have a positive impact on other properties, such as creep resistance, fatigue resistance, and oxidation resistance.

Data Availability Statement:
The data that support the findings of this study are available from the corresponding author upon reasonable request.

Conflicts of Interest:
The authors declare that they have no known competing financial interest or personal relationships that could have influenced the work reported in this paper.