Creep anisotropy reduction and improvement via post-heat treatment in yttrium-added Hastelloy-X fabricated by laser powder bed fusion

The development of laser powder bed fusion (LPBF) has made it possible to produce complex three-dimensional components for high-temperature applications. The LPBF process needs to be refined to address several key factors, such as high-temperature elongation, microstructure heterogeneity, and mechanical anisotropy. Hastelloy-X Ni-based superalloy was used to illustrate these issues in this study. First, 0.046 wt.% of yttrium (Y) was added to Hastelloy-X (HX-y) to prevent grain boundary embrittlement. The second step involves two kinds of post-heat treatments; i) at 850 °C for 2 h (DA) for carbides and ii) solution heat treatment at 1240 °C for 8 h upon aging at 850 °C for 2 h (HTA). The creep properties of the samples were compared at 900 °C/ 80 MPa to understand the effect of Y addition and post heat treatments. The HX-y specimen was strengthened by solid solution and dispersion of Y-rich oxides, together with stabilization of oxygen-based contamination at the grain boundaries. The DA and HTA HX-y specimens had better creep properties than the HX specimens. The HX-y specimen showed superior creep properties to the HX specimen due to the presence of carbides M6C and Cr23C6 inside grains and at grain boundaries. However, carbides remained stable even at high temperatures within grains and at grain boundaries. Nevertheless, the HTA HX-y specimen exhibited superior isotropic creep properties. As a result of grain boundary pinning, serrated grain boundaries prevented grain boundary sliding. In contrast, HX specimens exhibited poor creep properties. This study confirmed that the optimal addition of Y together with the optimization of post-heat treatment drastically enhances the creep properties of materials fabricated by the LPBF process.


Introduction
Rapid prototyping, also known as additive manufacturing (AM), enables new manufacturing technologies to make complex objects. Many industries are shifting from conventional manufacturing processes to AM. This process provides a means of creating models and printing parts, thus cutting down the amount of time, money, and human interaction involved in the product development cycle [1][2][3][4][5][6]. In contrast with subtractive manufacturing techniques such as machining, additive manufacturing refers to joining materials to make objects from 3D models layer by layer. By using additive manufacturing, intricate and complex geometries can be produced with minimal post-processing, using tailored materials such as plastics and metals [7][8][9].
Therefore, AM is a tool that provides designers and engineers with greater creative freedom and allows them to construct unique products in small quantities economically [10][11][12][13]. As a raw material for 3D printing, the powder is used in two main processes: selective laser melting (SLM) and electron beam melting (EBM) [14,15]. The SLM process is often used to melt powder material to make components. As each layer is built, the bed is lowered, and a fresh powder layer is applied. A rotating roller is then used to spread the powder evenly. An unsintered powder supports a sintered part while it forms the part [16,17]. After the build is complete, unsintered materials can be cleaned and recycled [18][19][20]. SLM has many advantages over other processes, such as a more refined microstructure, a smooth surface finish, and minimal post-machining. The SLM is also primarily used in metal additive manufacturing due to its high accuracy and surface quality [21][22][23]. However, SLM is still comparatively slow and suffers from nonuniform distributions of thermal fields, leading to heterogeneities in microstructure and random dispersion of defects like cracks and porosities [24][25][26]. Hastelloy-X is a Ni-based superalloy composed primarily of M 6 C, M 23 C 6 , σ, and μ phases. It is without γ′-precipitates and features good weldability [27]. It is used mainly for high-temperature aircraft engine combustor zone components. Much research has been done so far on Hastelloy-X using AM processing. Tomus et al. examined the effects of process parameters on microstructures and concluded that lower scan speeds increased density in SLM Hastelloy-X [28]. The influences of alloying elements on hot cracking, mainly Mn, Si, and C, have also been studied. Lower levels of these elements reduced hot crack formation during the SLM process [25].
Hana et al. reported the formation of hot cracks along the high-angle grain boundaries and interdendritic regions due to the interdendritic liquid pressure drop between the dendrite tip and root [26,29,30]. Anisotropic mechanical properties are general because of complex microstructures and crack formations in AM alloys [31][32][33][34]. Because of these critical issues, it is paramount for AM to obtain isotropic mechanical properties. We took this as a challenge to get isotropic creep properties by eliminating microstructural heterogeneity and cracks through post-heat treatment and the modification of composition. Yttrium (Y) was added to Hastelloy-X to enhance its mechanical properties. Banoth et al. have studied the effects of Y addition on alloys in detail [9,24,35].
Previously we produced crack-free Hastelloy-X by adding smaller amounts of Y and Si than in earlier studies [35]. The study involved heat treatment at 1177 °C for 2 h in the air, followed by creep and tensile tests at 900 °C. The results showed that the Y-added specimen exhibited better mechanical properties than a Y-free specimen produced under the same conditions. Primarily due to the carbides and yttria formation in the grain interiors and at grain boundaries. In the SLMed Y-added and Y-free specimens, the standard heat treatment did not remove the microstructural heterogeneity. In the present study, heat treatment at a higher temperature was proposed to enhance creep properties by reducing microstructural heterogeneity in Y-added and Y-free SLMed specimens. Heat treatments were selected from the Hastelloy-X TTT diagram [36]. In the present study, we aimed to determine how heat treatment conditions affected the creep properties of Y-added Hastelloy-X processed by an SLM. To determine the isotropic creep properties, tests were conducted perpendicular to and parallel to the direction of the building. Creep tests were conducted at 900 °C and 80 MPa. Table 1 presents the nominal chemical compositions of the two Hastelloy-X specimens: one without yttrium (HX) and the other with 0.046 wt.% yttrium (HX-y). We fabricated both alloys using the EOS M290 SLM machine (EOS, Krailling, Germany) with 45 mm 3 cube blocks in a protective Ar atmosphere (Fig. 1). Under different heat treatment conditions, yttrium's effects on microstructure and creep properties were investigated. The heat treatment conditions were chosen from the typical Hastelloy-X TTT diagram ( Fig. 2) [36]. The specimens performed aging heat treatment at 850 °C for 2 h in air, then air cooled (AC). To eliminate microstructural heterogeneity, a solution heat treatment was carried out at 1240 °C for 8 h, followed by aging at 850 °C for 2 h (HTA). The heat treatment conditions and specimen abbreviations are indicated in Table 2. The specimens were enclosed in vacuum-sealed quartz tubes to prevent oxidation at higher temperatures. The DA and HTA heat treatment conditions that we designed to eliminate the microstructural heterogeneity and their effects on creep properties.

Creep test
Creep rupture tests were performed using a temperature control unit (Toshin Kogyo, Tama-city, and Tokyo, Japan). Each condition was tested along normal to building direction (horizontal specimen) and along building direction (vertical specimen) at 900 °C with 80 MPa rupture stress. Cubes (45 mm 3 ) were built and then cut into several 3.1 mm-thick plates, from which the test specimens were cut by using an electro-discharge machine (EDM). The dimensions of the creep specimen gauge section are 19.6 × 2.8 × 3.0 mm.

Microstructural observation
The optical microscope (OM; Olympus, Tokyo, Japan) and an EDS (AMETEX 9424; EDAX, Tokyo, Japan) along with an FE-SEM (JSM-7200F; JEOL, Tokyo, Japan) were used to observe the microstructure. An accelerated voltage of 20 kV was used for the FESEM and EBSD analyses. Electron backscatter diffraction calculated inverse poles and pole figures

Grain size and recrystallization fraction calculations
A total of four IPF maps were taken at four different locations (× 100 magnification) to measure the average grain size and grain boundary fractions. Grain misorientation angles greater than 15° were considered grains. TSL software calculated the average grain size and the fraction of recrystallized grain boundaries. The grain boundaries were classified into three categories based on the grain boundary misorientation angle. The first category, grain boundaries with misorientation angles of 0-15°, are referred to as lowangle grain boundaries (LAGBs) and are shown in red on the grain boundary map. The second category, grain boundaries with misorientation angles of 15-35°, are shown in green as high-angle grain boundaries (HAGBs). The third category, grain boundaries with misorientation angles of 35-65°, are termed migrated high-angle grain boundaries (MHAGBs) [37]. Figure 3 shows optical microscopy images of HX and HX-y specimens. No cracks were observed in either specimen, though porosity was relatively lower. Epitaxial dendritic and interdendritic structures were observed in HX and HX-y specimens (Fig. 3c, d). Moreover, there was a dark phase formation along the grain boundaries (GB) of the HX-y specimen; this would be Y-rich oxides, as the HX-y specimen contained Y (Fig. 3e) [35].

Heat-treated specimens
Before the creep test, HX and HX-y DA vertical specimens were analyzed using FESEM and EDS ( Fig. 4a-c, d-f, respectively). After DA heat treatment, the grain morphology was similar to that of the as-built specimen in both specimens, but some cell structures were dissolved (Fig. 4a, d). However, white phases formed within the grains and at the grain boundaries in HX and HX-y DA specimens (Fig. 4c, f). The EDS analysis revealed the formation of Mo 6 C and Cr 23 C 6 carbides inside the grains in the HX DA specimen (Fig. 5a). In the HX-y DA specimen, Y-rich and Mo 6 C phases formed at the grain boundaries and interior of the grains (Fig. 5b, c). Cr 23 C 6 carbides are also found at the grain boundaries (Fig. 5c). As a result of the HTA heat treatment, both HX and HX-y specimens exhibited equiaxed grain morphologies (Fig. 6). The morphology of grain boundaries differed significantly between the HX-y HTA specimen and the HX HTA specimen. The HX-y HTA specimen exhibited serrated grain boundaries and these were uniform throughout the specimen (Fig. 6e). In addition, bright and dark phases were prominent along the grain boundaries and within the grains (Fig. 6f). They were rich in Mo, Cr, and Y (Fig. 6g, f). The HX HTA specimen showed small numbers of tiny bright and dark phases at the grain boundaries; however, this specimen demonstrated straight grain boundaries due to a lack of precipitates forming at the grain boundaries ( Fig. 6c).

Creep properties
The creep curves for specimens HX and HX-y are shown in Fig. 7a at as-built and DA conditions. In both vertical and horizontal directions, the creep life of the HX asbuilt specimen was inferior to that of the HX-y specimen (Fig. 7, Table 3). In comparison with the HX-y DA specimen, the HX vertical specimen had poor creep properties after DA heat treatment. Compared to the HX DA specimen, the HX-y vertical specimen had a significantly longer creep life (98 h), 7.66 times as long as the HX DA. As shown in Table 3, the rupture elongation for HX-y DA vertical specimen (8.11%) was 2.42 times greater than that for HX DA vertical specimen. Figure 7b presents the creep rate versus time curves in the vertical direction. The creep rates were the same for both as-built and DA HX-y specimens. In contrast, the as-built and DA HX specimens showed a higher creep rate (Fig. 7b, Table 3). The creep curves of the horizontal specimens HX and HX-y at both as-built and DA conditions are shown in Fig. 7c and Table 3. The creep life and rupture elongation of horizontal specimens were both reduced to 1.72 h and 0.64, respectively. In contrast, the as-built horizontal HX-y specimens had a longer creep life (32.7 h) and greater rupture elongation (5.48%). As illustrated in Fig. 7c, Table 3, the DA horizontal HX-y specimens also showed better creep properties than the HX DA horizontal specimens. As shown in Fig. 8a, b, both horizontal and vertical HTA specimens for HX-y had the same creep life, rupture elongation, and creep rate after HTA heat treatment (Table 3). Despite being inferior to HX-y HTA specimens ( Fig. 8 and Table 3), the HX HTA vertical and horizontal specimens also exhibited isotropic creep properties. Figure 9 shows the fracture surface of a vertical specimen. It was observed that cracks nucleated and propagated along grain boundaries and interdendritic regions in HX and HX-y DA specimens, similar to the as-built HX-y specimens, corresponding to the cleavage fracture. Conversely, intergranular fractures predominated in HX DA and HX-y DA specimens (Fig. 9a HX DA and b HX-y DA). As shown in Fig. 9c, d, the fracture surfaces of the HX HTA specimen are smooth, while those of the HX-y HTA specimen are coarse (Fig. 9e, f). Figure 10 shows a top view of the fractured surface at the grain boundary. At the grain boundaries, small, round, and uniform white phases can be seen in the HX-y HTA specimen (Fig. 10a, b). Figure 10c, d demonstrate that Mo-and Y-rich phases have formed based on the elemental mapping and line analysis results from EDS. In contrast, no phases were observed on the fractured surface of the HX HTA specimen (Fig. 9c, d).  The vertical specimens of HX and HX-y after the creep test are shown in Fig. 11. Both HX and HX-y crept specimens showed crack nucleation and propagation paths along grain boundaries. In the HX DA specimen, white phases were observed along the grain boundaries and inside the grain (Fig. 11c) and were rich in Mo, W, and C. Compared to those in the HX DA crept specimen (Fig. 11c), the grain boundary carbides are much larger in the HX-y DA specimen (Fig. 11f). The EDS point analysis revealed that location 1 contained significant amounts of Mo, C, and to a lesser extent, W and Y. These must have been (Mo, W) 6 C carbides and Y-rich oxides (Table 4). Cr 23 C 6 carbides also formed at the grain boundary. The presence of Y in the matrix phase (γ) was also clear.

Fractography and crept specimen microstructures
The side view of the fractured surfaces of the HX and HX-y HTA vertical specimens revealed no difference between crack initiation and propagation (Fig. 12a-c and d-f); this resulted in the nucleation of cracks and the propagation of those cracks along grain boundaries (Fig. 12c, f). However, one key difference was observed regarding the formation of carbides: they were confined primarily to the grain boundaries in the HX HTA specimen. Their size was smaller than that of the HX-y HTA specimen (Fig. 12b). Conversely, carbide formation at grain boundaries and inside the grain was more apparent in the HX-y HTA specimen. In HX-y HTA specimens, grain boundary carbides were larger and discrete (Fig. 12e). The results of the EDS point analysis on both the HX and HX-y HTA specimens are presented in Table 5. The HX HTA specimen at locations 1 and 2 showed a pattern of (Mo, W) 6 C and Cr 23 C 6 carbides (Fig. 12c, Table 5). At the same time, the results for the HX-y HTA specimen were also the same as those for the HX HTA specimen, except for the Y-rich phase formation and a solid solution of Y in the matrix phase (γ) (Fig. 12f, Table 5).

Recrystallization via post-heat treatment and its effect on creep properties
An additive manufacturing part is typically subjected to a complex thermal cycle that involves the extraction of directional heat, the repetition of melting, and the rapid solidification of the part [38,39]. The process would result in microstructures that are anisotropic and non-homogeneous, which would result in parts that are intrinsically different from those that are produced through conventional methods. Thus, parts created by metal additive manufacturing may possess anisotropic and heterogeneous properties [10,31]. Due to defects such as pores and rough surfaces, as well as the absence of fusion layers, metal AM parts can also exhibit anisotropic and heterogeneous properties [40,41]. This study did not consider the influence of defects such as pores, micro-cracks, and surface roughness on the anisotropy and heterogeneity properties of SLMed Hastelloy-X. Moreover, there are two ways to alter the microstructural heterogeneity of material after AM to achieve isotropic mechanical properties. First, hot isostatic pressing (HIP) is used to compress materials by applying a high temperature of several hundred to 2000 °C as well as an isostatic pressure of several tens to 200 MPa simultaneously in an inert gas environment. However, the main disadvantage of the hot isostatic pressing process is the cost of the tooling. This makes them only applicable to special applications like aerospace [42]. Secondly, heat treatment at a higher temperature for an extended period can modify the microstructure of an AM material. As far as cost and machinery are concerned, hightemperature furnaces are readily available for industrial and academic use. The present study used the second method to achieve homogeneous microstructures. In this study, two main heat treatment methods were applied based on the Hastelloy-X TTT diagram. The SLM HX and HX-y specimens are heat treated at 850 °C in the air for 2 h, then cooled in the air (DA). Its primary purpose is to examine the effects of carbides on creep properties. A second heat treatment involved solution heat treatment at 1240 °C for 8 h, followed by aging at 850 °C for 2 h in Ar gas-filled capsules, followed by air cooling (HTA). It is intended to eliminate microstructural heterogeneity both in HX and HX-y specimens. The pole figures (PF) of the HX and HX-y as-built specimens along the building and vertical directions demonstrated that both specimens prevailed in columnar grains with < 001 > crystallographic orientation (Figs. 13a, c and  14a, c). Figure 14b, d shows the grain boundary misorientation angles of the as-built HX and HX-y specimens. Both specimens had almost equal grain boundary misorientation angles in the as-built conditions; that is, both specimens had maximum grain boundaries in the range of 0-10° angles (Fig. 14b, d). Compared to the HX as-built specimen, the HX-y as-built specimen exhibited good creep properties in vertical and horizontal directions (Fig. 7). It is possible that the Y addition may have caused the solid solution strengthening and carbide formation inside the grains and at the grain boundaries. A detailed explanation is given in Sect. 4.2.
Because the DA heat treatment is designed to get carbides in the present study, both HX specimens and HX-y specimens maintained columnar grain morphology after the DA heat treatment (Fig. 13e, g). Moreover, the aging heat treatment at 850 °C caused partial recrystallization in HX and HX-y DA specimens. As a result, the dislocation density decreased. At the same time, the MHAGBs were increased (Fig. 14f, h and Table 6). As a result, the HX-y specimen had a creep life of 98 h along the vertical direction, which is 7.66 times longer than the creep life of the HX DA vertical specimen. Compared to the HX DA vertical specimen, the rupture-elongation of the HX-y DA vertical specimen was 8.11%, which was higher than the HX DA vertical specimen However, after the DA heat treatment, microstructural heterogeneity persisted, causing anisotropic creep properties in both HX and HX-y DA specimens (Fig. 7a, c). Since the DA heat treatment purpose is to precipitate the carbides in γ-matrix and its effect on the creep properties.
To eliminate microstructural anisotropy, the solution heat treatment was carried out at 1240 °C for 8 h followed by aging at 850 °C for 2 h. This led to complete recrystallization in the HX and HX-y HTA specimens (Fig. 13i, k). At the same time, the grain morphologies changed to equiaxed. The SLM process is known to cause residual stresses in the materials due to the dynamic heating and cooling, which results in numerous dislocations in the material that were eradicated after HTA heat treatment [43,44]. Moreover, MHAGBs and grain boundary misorientation angles were increased (Fig. 13j, l and j, l). The grain aspect ratio and grain boundary misorientation angles also increased ( Table 6). This indicates that the grain shape became equiaxed and more grain boundaries had higher misorientation angles (Figs. 13l and 14l). The change in microstructure after HTA heat treatment can provide a clear interpretation of the creep properties of the HX and HX-y specimens. The creep properties of the vertical and horizontal specimens of the HX-y HTA are illustrated in Fig. 8a and Table 3. In both vertical and horizontal directions, the creep properties were approximately identical, indicating isotropic creep. Both vertical and horizontal HX-y HTA specimens exhibited the same creep rate (Fig. 8b, Table 3). Alternatively, the HX HTA specimens exhibit similar creep properties in both vertical and horizontal directions but are inferior compared to HX-y HTA specimens (Fig. 8).

Effects of Y and carbides on creep behavior
Hastelloy-X is a Ni-based superalloy that is strengthened by solid solution and carbides. The main elements in strengthening solid solution are Mo, Cr, and W, and carbides are strengthened mainly by M 6 C and M 23 C 6 . The M 6 C carbides form grain interiors at grain boundaries, whereas the M 23 C 6 carbide is formed primarily at grain boundaries. In addition, there are detrimental phases called σ and μ, which form after long exposures to high temperatures. In the present study, the aging heat treatment (DA) enhanced the creep life along the vertical direction in the HX-y specimen from the asbuilt specimen (Fig. 7a). Two factors may have caused this better creep resistance. One is the solid solution strengthening effect due to the addition of Y (0.046 wt.%), so solubility in the matrix is possible. In addition, the atomic radius of Y (2.27 Å) is larger than that of Ni (1.62 Å). So Y in the matrix has a good effect on solid solution strengthening. Li et al. showed a positive effect by the addition of optimum Y in the Ni-Mo-Cr-Fe Ni-based superalloy [45]. The second factor that may explain the enhancement of the creep properties is carbide formation. The formation of carbides was observed in both HX DA and HX-y DA specimens (Fig. 4). The EDS elemental mapping of the HX-y DA specimen illustrates the formation of carbides such as Mo 6 C and Cr 23 C 6 (Fig. 4c). It is also evident that M 6 C carbides in the HX-y DA specimen after the creep test were much larger than those in the HX DA specimen (Fig. 11d-f, Table 4 and 7). Y addition was found to promote carbide growth.
After the HTA heat treatment, the carbides in the HX-y specimen are larger at the grain boundaries and inside the grains (Fig. 6d-f and Table 7). The carbides that formed in the HX-y specimen may have been a consequence of the yttrium carbide (YC) that formed during the SLM process   [35]. The YC can serve as a nucleation site for carbides and oxides within grains [24,46]. As a result of high-temperature exposure, carbide precipitation may also result from the segregation of the carbon in areas with a high density of dislocations, stacking faults, or grain boundaries [47]. During solidification, the larger atomic elements are discharged to the liquid phase and segregate at the interdendritic regions (L → γ + M 6 C (M stands for Mo, W, Cr, and Y)) [45]. Because the Y atomic radii is larger, it will resist Mo and Cr diffusion from the interdendritic, matrix, and grain boundary regions [48,49]. At the same time, carbon is an interstitial atom; it readily diffuses to grain boundaries due to long-range diffusion, which may have caused the formation of carbides during the solution heat treatment and aging in the HX-y specimen. Thermodynamic calculations were done to analyze the Y effect on carbides. We considered a nonequilibrium or Scheil model to study the carbide stability in the Y-added alloy. We used Pandat 2021 software and the PanNi 2020 database for thermodynamic analysis. This analysis supports the experimental results regarding the formation and stability of carbides in the HX-y specimen (Fig. 15). In the HX specimen, the M 6 C carbide, which is the major and eutectic phase in the alloy, is stable at temperatures above 1300 °C (Fig. 15a, b). In comparison, in the Y-added HX-y specimen, the M 6 C carbide is stable from 1150 °C to 1350 °C (Fig. 15c, d). In the present study, the HX-y HTA heat treatment was conducted at 1240 °C, where M 6 C carbide would be stable, but it would not be stable in the case of the HX HTA specimen. Moreover, the grain boundaries were serrated in the HX-y HTA specimen; serrated boundaries were observed throughout the specimen (Fig. 6e, f). The serrations formed due to the presence of hard phases rich in Mo, Cr, C, and Y oxides at the boundaries, and these phases would have pinned the grain boundaries during the solution heat treatment at 1240 °C (Fig. 6e, f). Serrated grain morphologies enhance creep properties by inhibiting grain boundary sliding [50]. There was a distinct difference between the fracture surfaces of the HX and HX-y HTA specimens: the HX HTA specimens were smooth, whereas the HX-y HTA specimens were coarse (Fig. 9c, d). A grain boundary may cause a coarse surface with many phases rich in Mo and Y (Fig. 10c, d). In the HX-y HTA specimen, grain boundary precipitates contribute to better grain boundary sliding resistance so that the creep properties are superior to those of the HX HTA specimen (Fig. 8). In addition, it was found that M 6 C carbides are much larger in the HX-y HTA specimen after the creep test than in the HX HTA specimen (Fig. 12d-f, Table 4). In both the HX and HX-y alloys, the carbides in the grains would be nucleation sites of creep void formation (Fig. 12c, d). Cao of Y addition on carbide precipitation in the K4169 superalloy. By increasing the yttrium content from 0.005 wt.% to 0.033 wt.%, the MC carbides' morphology and size are greatly affected, since the yttrium alters carbide diffusion during solidification [51].
An excessive amount of oxygen in the starting material caused grain-boundary embrittlement, which adversely impacted the AM process. To eliminate this problem in the AM process, Hastelloy-X was enriched with yttrium. Furthermore, the AM process results in anisotropic microstructures since melting and solidification are localized. To achieve isotropic creep properties, we used a unique heat treatment to eliminate microstructural heterogeneity and examined how the treatment affected crack-free Y-added Hastelloy-X manufactured by an LPBF process. The purpose of this study is to offer some insight into the manufacturing of complex structural components that have superior mechanical properties and can be used in high-temperature applications by additive manufacturing.

Conclusions
We identified typical microstructures of AM alloys. As a result of the LPBF process, the HX and HX-y alloys exhibited an isotropic textured microstructure with columnar grain morphology. To eliminate anisotropy, unique post-heat treatments were applied. Based on this study, the following conclusions were drawn.
1. The HX-y specimen exhibited enhanced creep properties following the DA and HTA heat treatments compared with the HX specimen. 2. After HTA heat treatment, the HX-y specimen exhibited complete recrystallization. It thus achieved isotropic creep properties. As a result, the creep life and rate were equal in the vertical and horizontal directions. 3. Because of the formation of carbides, serrated grain boundaries, and solid solution strengthening due to the Y addition, the HX-y HTA specimen's creep properties were superior to those of the HX HTA specimen.