E ﬀ ects of Yttrium on Microstructure Stability and Tensile Properties of China Low Activation Martensitic Steel

: This study investigated the microstructural stability and mechanical properties exhibited by China low activation martensitic (CLAM) steels with di ﬀ erent yttrium (Y) contents over 3000 h of aging at 550 ◦ C. Scanning electron microscopy, electron backscatter di ﬀ raction analysis, and transmission electron microscopy were employed to investigate the microstructural evolution of the steels. Results indicated that grain boundary migration was slow and the Laves phase precipitation was delayed in Y-containing steels. Grain boundaries at di ﬀ erent angles in 0Y and 6Y CLAM steels were signiﬁcantly a ﬀ ected, and those in 36Y and 71Y alloys exhibited negligible changes during the long-term thermal aging. Moreover, Y contents had appreciable e ﬀ ects on the strength and toughness of the aged steels. The stable microstructure of Y-containing CLAM alloys is responsible for improved strength and impact toughness during aging.


Introduction
Reduced activation ferrite/martensitic (RAFM) steels with 9 wt.% Cr are the most promising structural materials for use in future fusion reactors because of their good thermomechanical properties and creep resistance at high temperatures [1][2][3]. RAFM steels with guaranteed low activation have been developed by replacing Nb and Mo in conventional 9Cr-FM steels with low-activated Ta and W [4,5]. Numerous countries have developed their own RAFM steels [6,7], such as Eurofer97 (Europe), F82H (Japan), 9Cr-2WVTa (USA), and China low activation martensitic steel (CLAM) (China). CLAM, 9Cr-1.5W-0.2V-0.15Ta, was developed by the Institute of Plasma Physics, Chinese Academy of Sciences (ASIPP), and other institutes and universities [8]. The smelting process [9], mechanical properties [10], compatibility with liquid LiPb [11], and further research and development [12] of CLAM have been evaluated by the ASIPP.
Rare earth (RE) elements have become an important additive in the metallurgical industry, being used as deoxidizers, desulfurizers, modifiers, and alloying elements because of their 4f electron characteristic, alternation valence, large atomic size, and enhanced chemical activity [13]. Numerous studies have focused on the RE element cerium [14]. Another RE element, yttrium (Y), shares similarities with cerium while also possessing unique characteristics [15]. In addition, Y 2 O 3 and Y are widely used in strengthening phases in oxide dispersion-strengthened (ODS) alloys [16]. The addition of 0.2 wt.% Y Samples for testing were spark-machined from aged plates along the rolling direction. A layer with a thickness of more than 2 mm was removed from the surface of each specimen to eliminate the influence of oxidation on thermal aging properties. Specimens for optical microscopy (Leica, Solms, Germany) and scanning electron microscopy (SEM) (ZEISS, Jena, Germany) were etched with vilella reagent (100 mL of alcohol + 5 mL of picric acid + 1 g of muriatic acid). The average prior grain sizes (APGZ) of the steels were measured using the linear intercept method. Transmission electron microscope (TEM) (FEI, Hillsboro, OR, USA) specimens were first mechanically polished to be 30-50 µm thick. Then, a thin disk with a diameter of 3 mm was subjected to twin-jet polishing with an electrolyte solution (95% acetic acid + 5% perchloric acid) at room temperature (RT). Rod specimens with a diameter of 5.0 mm and a length of 25.0 mm were used to test strength at RT. Dog-bone-shaped flat specimens (6.0 mm × 2.0 mm × 30 mm) were used for testing at 550 • C, 600 • C, and 650 • C. Tensile tests were performed with a cross-head speed of 2 mm/min. Full-sized V-notch Charpy specimens (10 mm × 10 mm × 55 mm) were prepared for impact tests over a temperature range of −130 to 25 • C. All tensile tests were performed three times, and the results were averaged. The schematics of the specimens are shown in Figure 1.
with a diameter of 5.0 mm and a length of 25.0 mm were used to test strength at RT. Dog-bone-shaped flat specimens (6.0 mm × 2.0 mm × 30 mm) were used for testing at 550 °C, 600 °C, and 650 °C. Tensile tests were performed with a cross-head speed of 2 mm/min. Full-sized V-notch Charpy specimens (10 mm × 10 mm × 55 mm) were prepared for impact tests over a temperature range of −130 to 25 °C. All tensile tests were performed three times, and the results were averaged. The schematics of the specimens are shown in Figure 1.

Grain Structural Characteristics
The microstructures of the CLAM alloys after long-term aging are presented in Figure 2. The original heat-treated matrix [19] remained unchanged after aging at 550 • C for 3000 h. Martensitic lath was present in all samples. Table 2 shows the APGZ of the aged samples measured through the linear intercept method. Grains coarsened during the 3000 h of aging but were refined through the hindrance of grain boundaries by Y-containing inclusions, as previously confirmed [19]. The lower coarsening rate of the CLAM steel with 0.071 wt.% Y (11.6 → 13.1 µm) as compared to that of the reference base steel (0Y alloy) (14.7 → 16.9 µm) was attributed to the pinning of the Y inclusions [19]. Surface energy is the driving force of grain boundary migration [20]. Thus, grain coarsening during the 3000 h of aging was a spontaneous process.

Grain Structural Characteristics
The microstructures of the CLAM alloys after long-term aging are presented in Figure 2. The original heat-treated matrix [19] remained unchanged after aging at 550 °C for 3000 h. Martensitic lath was present in all samples. Table 2 shows the APGZ of the aged samples measured through the linear intercept method. Grains coarsened during the 3000 h of aging but were refined through the hindrance of grain boundaries by Y-containing inclusions, as previously confirmed [19]. The lower coarsening rate of the CLAM steel with 0.071 wt.% Y (11.6 → 13.1 μm) as compared to that of the reference base steel (0Y alloy) (14.7 → 16.9 μm) was attributed to the pinning of the Y inclusions [19]. Surface energy is the driving force of grain boundary migration [20]. Thus, grain coarsening during the 3000 h of aging was a spontaneous process.

Precipitations
Additional microstructural details were investigated through back-scattered electron imaging (BSE). New kinds of phases were observed on grain boundaries in the 0Y alloy that had been aged for 1500 h. These phases are marked with black arrows in Figure 3a. The inclusion of elemental Fe, Cr, and W in the phases (Figure 3i) was indicative of the Laves phase ((Fe, Cr)2W) observed in other

Precipitations
Additional microstructural details were investigated through back-scattered electron imaging (BSE). New kinds of phases were observed on grain boundaries in the 0Y alloy that had been aged for  Figure 3a. The inclusion of elemental Fe, Cr, and W in the phases (Figure 3i) was indicative of the Laves phase ((Fe, Cr) 2 W) observed in other aged RAFM steels [21]. Laves phases were almost absent from Y-containing CLAM steels aged at 550 • C for 1500 h (Figure 3b-d). The atomic radius of Y, an RE element, is approximately 40% larger than that of Fe atoms. Therefore, the tendency of Y to occupy vacancies, dislocations, phase interfaces, grain boundaries, and other defects upon addition to steel reduces the free energy of the system [22,23]. The segregation of RE elements at grain boundaries would hinder the segregation of elemental W, Cr, and C at grain boundaries and affect the nucleation and precipitation of second phases by considerably affecting the solubility of other alloying elements at grain boundaries [22,23]. The solubility of RE elements would increase when additional RE elements are added, while the carbide nucleation would intensify. The presence of residual Y-rich inclusions in the 71Y alloy would destroy the continuity of the matrix and degrade the mechanical property of the alloy [19,24], as shown in Figure 3d. Several Laves phases appeared in the Y-added CLAM steels and the number of Laves phases in the 0Y alloy increased (Figure 3e-h) when the exposure time was extended to 3000 h. The solubility of RE elements would increase when additional RE elements are added, while the carbide nucleation would intensify. The presence of residual Y-rich inclusions in the 71Y alloy would destroy the continuity of the matrix and degrade the mechanical property of the alloy [19,24], as shown in Figure 3d. Several Laves phases appeared in the Y-added CLAM steels and the number of Laves phases in the 0Y alloy increased (Figure 3e-h) when the exposure time was extended to 3000 h.
Blocky yttrium-rich inclusions The details of the microstructures of the CLAM steels were revealed through TEM ( Figure 4). Martensitic lath boundaries were observed in the aged samples. This observation corresponds with the observation inferred from SEM images ( Figure 3). The tempered martensitic microstructure did not show any changes except for lath coarsening during aging at 550 °C. The widths of the lath martensite in the aged alloys are shown in Table 3. Block widths increased by 24.1% (0.29 → 0.36 μm) in the 36Y alloy, by 70.3% (0.37 → 0.63 μm) in the 0Y alloy, by 28.6% (0.28 → 0.36 μm) in the 6Y alloy, and by 36.4% (0.33 → 0.47 μm) in the 71Y alloy. Grain boundaries gradually evolved during the 3000 h of aging. Lath coarsening is a process of dislocation motion and annihilation, and the prolongation of aging time could increase this [25]. The boundary of the martensite lath could have coarsened with time, but significant evidence for the evolution of martensitic lath into sub-grains was not obtained because of the low aging temperature [26]. Meanwhile, precipitates on grain boundaries and inside grains increased after aging. Numerous inhomogeneous spherical, rod, and ellipsoid M23C6 carbides could be found on prior austenite grain and martensitic lath boundaries. The nucleation and growth of M23C6 carbides were controlled by interfacial energy and the segregation of Cr and C elements. The reduction in the interfacial energy of grain boundaries by RE elements that had segregated at boundaries could inhibit the segregation of elemental Cr. Meanwhile, the activity of elemental C could be reduced, and the solution of elemental C could be promoted by RE elements. The details of the microstructures of the CLAM steels were revealed through TEM (Figure 4). Martensitic lath boundaries were observed in the aged samples. This observation corresponds with the observation inferred from SEM images (Figure 3). The tempered martensitic microstructure did not show any changes except for lath coarsening during aging at 550 • C. The widths of the lath martensite in the aged alloys are shown in Table 3. Block widths increased by 24.1% (0.29 → 0.36 µm) in the 36Y alloy, by 70.3% (0.37 → 0.63 µm) in the 0Y alloy, by 28.6% (0.28 → 0.36 µm) in the 6Y alloy, and by 36.4% (0.33 → 0.47 µm) in the 71Y alloy. Grain boundaries gradually evolved during the 3000 h of aging. Lath coarsening is a process of dislocation motion and annihilation, and the prolongation of aging time could increase this [25]. The boundary of the martensite lath could have coarsened with time, but significant evidence for the evolution of martensitic lath into sub-grains was not obtained because of the low aging temperature [26]. Meanwhile, precipitates on grain boundaries and inside grains increased after aging. Numerous inhomogeneous spherical, rod, and ellipsoid M 23 C 6 carbides could be found on prior austenite grain and martensitic lath boundaries. The nucleation and growth of M 23 C 6 carbides were controlled by interfacial energy and the segregation of Cr and C elements. The reduction in the interfacial energy of grain boundaries by RE elements that had segregated at boundaries could inhibit the segregation of elemental Cr. Meanwhile, the activity of elemental C could be reduced, and the solution of elemental C could be promoted by RE elements. The precipitation of M 23 C 6 carbides were inhibited by elemental Y as a result. The carbide content of CLAM steels with Y was lower than that of 0Y alloy, as shown in Figure 4. However, a large number of carbides were observed on the grain boundaries of alloys with Y after 3000 h of exposure at 550 • C.    Figure 5 shows the distribution of the misorientation angles in CLAM steels during aging. The angles were mainly less than 10° and ranged from 50 to 60°. The presence of a maximum peak that was less than 5° indicated that sub-grain boundaries and lath boundaries were the main boundaries in the CLAM steels. The same results were reported for CLAM steel that had been aged at 550 °C for 10,000 h [27], and 9Cr-ODS steel that had been aged at 700 °C for 10,000 h [28]. In the present study, grain boundary angles that were less than 10° were defined as low-angle grain boundaries (LABs) and those that were greater than 10° were defined as high-angle grain boundaries (HABs). Long-term thermal aging had a drastic effect on the percentages of grain boundaries with different angles, particularly 4-10° LABs and 50-60° HABs, in 0Y and 6Y CLAM steels. The misorientation angle distributions of 36Y and 71Y alloy exhibited negligible changes during aging. Table 4 shows the statistical results for misorientation angles in the alloys. Wang et al. [29] and Karthikeyan et al. [30] classified original austenite grain boundaries in martensitic steels as HABs, and sub-grain and martensite lath boundaries as LABs. The increase in the proportion of LABs indicated that sub-grains gradually formed during aging at 550 °C. The coarsening and growth of austenite grains would be responsible for the reduction in HABs in steel during aging and the reduction of the unit volume of HABs. The same result could be inferred from the TEM images ( Figure 4) and the statistical results for average grain sizes shown in Table 2.   Figure 5 shows the distribution of the misorientation angles in CLAM steels during aging. The angles were mainly less than 10 • and ranged from 50 to 60 • . The presence of a maximum peak that was less than 5 • indicated that sub-grain boundaries and lath boundaries were the main boundaries in the CLAM steels. The same results were reported for CLAM steel that had been aged at 550 • C for 10,000 h [27], and 9Cr-ODS steel that had been aged at 700 • C for 10,000 h [28]. In the present study, grain boundary angles that were less than 10 • were defined as low-angle grain boundaries (LABs) and those that were greater than 10 • were defined as high-angle grain boundaries (HABs). Long-term thermal aging had a drastic effect on the percentages of grain boundaries with different angles, particularly 4-10 • LABs and 50-60 • HABs, in 0Y and 6Y CLAM steels. The misorientation angle distributions of 36Y and 71Y alloy exhibited negligible changes during aging. Table 4 shows the statistical results for misorientation angles in the alloys. Wang et al. [29] and Karthikeyan et al. [30] classified original austenite grain boundaries in martensitic steels as HABs, and sub-grain and martensite lath boundaries as LABs. The increase in the proportion of LABs indicated that sub-grains gradually formed during aging at 550 • C. The coarsening and growth of austenite grains would be responsible for the reduction in HABs in steel during aging and the reduction of the unit volume of HABs. The same result could be inferred from the TEM images ( Figure 4) and the statistical results for average grain sizes shown in Table 2.

Tensile properties
The tensile properties of the alloys at room temperature are shown in Table 5. The yield strength (YS) of the alloys at 550 °C, 600 °C, and 650 °C is plotted in Figure 6. The YS and ultimate tensile strength of the 0Y alloy first increased and then decreased with the prolongation of aging time at RT, as shown in Table 5. The increment in YS could be attributed to the precipitation of Laves phases, and the decrement in YS could be attributed to the coarsening of carbides and grains. For the steels with Y, YS first decreased and then held the line with a small fluctuation. Grain coarsening dominated during the early aging stage, and carbide precipitation mainly accounted for the fluctuation observed later in the aging stage. The trends shown by the changes in the YS of the 0Y alloy with aging time remained consistent as the test temperature increased. A temperature inflection point (600 °C) existed for the 36Y and 71Y steels. The YS first increased and then remained consistent because of the reduction in sub-grains at 600 °C. In contrast to the low-temperature stage (<600 °C), the precipitation of the second phase dominated in this stage (≥600 °C). The inclusions in Y-CLAM alloys were mainly

Tensile Properties
The tensile properties of the alloys at room temperature are shown in Table 5. The yield strength (YS) of the alloys at 550 • C, 600 • C, and 650 • C is plotted in Figure 6. The YS and ultimate tensile strength of the 0Y alloy first increased and then decreased with the prolongation of aging time at RT, as shown in Table 5. The increment in YS could be attributed to the precipitation of Laves phases, and the decrement in YS could be attributed to the coarsening of carbides and grains. For the steels with Y, YS first decreased and then held the line with a small fluctuation. Grain coarsening dominated during the early aging stage, and carbide precipitation mainly accounted for the fluctuation observed later in the aging stage. The trends shown by the changes in the YS of the 0Y alloy with aging time remained consistent as the test temperature increased. A temperature inflection point (600 • C) existed for the 36Y and 71Y steels. The YS first increased and then remained consistent because of the reduction in sub-grains at 600 • C. In contrast to the low-temperature stage (<600 • C), the precipitation of the second phase dominated in this stage (≥600 • C). The inclusions in Y-CLAM alloys were mainly Fe-Cr-Y-Ta-S-O, which act as strengthening particles in the steel [18,19,24]. These results indicate that the addition of Y affected the precipitation of carbides.

Charpy Impact Properties
The absorbed impact energy values of the aged alloys are shown in Figure 7. The Boltzmann function ( , where x0 is the ductile-brittle transition temperature) was used to fit the impact curves to obtain the ductile-brittle transition temperature (DBTT). The DBTT of the steels are shown in the upper left-hand side of Figure 7. The impact curve exhibited two distinct regimes. The upper shelf regime (USR) ranged from −20 to 25 °C. The DBTT drastically increased and

Charpy Impact Properties
The absorbed impact energy values of the aged alloys are shown in Figure 7. The Boltzmann function (A KV = , where x 0 is the ductile-brittle transition temperature) was used to fit the impact curves to obtain the ductile-brittle transition temperature (DBTT). The DBTT of the steels are shown in the upper left-hand side of Figure 7. The impact curve exhibited two distinct regimes. The upper shelf regime (USR) ranged from −20 to 25 • C. The DBTT drastically increased and the USR decreased with aging. Cleavage crack propagation was retarded or deviated because of the high grain boundary area per unit volume attributed to small grain sizes. The 36Y alloy showed a grain size that could improve toughness under the same aging time. Crack propagation could be effectively retarded or deflected by HABs, such as martensitic lath boundaries. The grain size of the steels increased by  Table 2. The relationship between DBTT and grain size could be expressed as follows [27]: where α and β are material constants, and d eff is the effective grain size. This relationship indicates that DBTT increased as the effective grain size increased.

Conclusion
The microstructural stability and mechanical properties exhibited by CLAM steels with different Y contents over 3000 h of aging at 550 °C were investigated. The following conclusions were obtained: (1) Long-term thermal aging drastically affected the percentages of grain boundaries at different angles in 0Y and 6Y CLAM steels. This effect indicates that sub-grains formed in the alloys during aging for 3000 h.
(2) The addition of elemental Y could improve microstructural and carbide stability. The mechanical properties of 6Y and 36Y alloys were better than those of 0Y steel. However, residual blocky Y-rich inclusions in 71Y alloy could degrade the performance of the aged alloy.
(3) The precipitation of the Laves phase was the main reason for the reduction in USR and the increment in DBTT.