Graphical abstract

1. Introduction

Fe-ODS are promising structural materials for high-temperature applications because they combine a ferritic matrix with a stable dispersion of fine oxide-rich particles [1]. These particles can hinder dislocation motion, reduce grain-boundary mobility, and help maintain microstructural stability during thermal exposure [2, 3]. Unlike conventional ferritic steels, the performance of ODS steels is governed not only by matrix composition, but also by how processing controls grain structure, deformation substructure, and particle distribution [2, 4-6].

The mechanical response of Fe-based ODS steels is strongly linked to the coupled composition–processing–microstructure relationship. Mechanical alloying and hot consolidation introduce dispersed oxide-rich particles into the ferritic matrix, while subsequent thermo-mechanical processing can produce refined grains, crystallographic texture, low-angle boundaries, and local stored deformation energy. These features jointly control the balance between strength and ductility [2, 3, 5]. Therefore, tensile behavior in ODS steels cannot be interpreted from composition or processing parameters alone; it must be correlated with the matrix deformation structure and the expected oxide-dispersion-related strengthening effects.

In multi-component Fe–15Cr ODS, alloying additions add further complexity. Mo and W contribute to solid-solution strengthening and may help maintain strength at elevated temperature. Ti is commonly associated with the formation of fine Y–Ti–O-type oxide particles, while Zr can modify oxide chemistry and may promote complex Y–Zr–O or Y–Ti–Zr–O features [6-8]. In addition, carbon may participate in secondary carbide formation, which can further affect local deformation and grain-boundary behavior. These interactions make the heat-treatment response of multi-component Fe–15Cr ODS more complex than that of simpler Fe–Cr–Y₂O₃ or Fe–Cr–Ti–Y₂O₃ systems.

A key issue in hot-rolled ODS steels is the competition between recovery, recrystallization-like boundary evolution, and particle-assisted microstructural stabilization. Hot rolling introduces stored deformation energy through dislocations, subgrain boundaries, and local orientation gradients. During subsequent heat treatment, this stored energy can drive recovery and grain-boundary migration. At the same time, oxide-rich particles may restrict boundary motion, depending on their size, spacing, distribution, and thermal stability [6, 8, 9]. The final microstructure is therefore expected to reflect a balance between deformation recovery, grain growth, and particle-related stabilization.

Despite extensive studies on extruded, forged, and tube-processed ODS steels, the heat-treatment response of hot-rolled multi-component Fe–15Cr ODS steels remains insufficiently clarified. In particular, it is still unclear how Mo–W–Zr–Ti–C co-alloying affects recovery and grain-boundary evolution after hot rolling, and how the resulting EBSD-scale deformation structure relates to tensile behavior. This gap is important because post-rolling heat treatment may improve ductility by reducing local misorientation and stored deformation, but it may also reduce strength through grain coarsening and substructure relaxation.

In this study, a hot-rolled Fe–15Cr ODS steel was investigated before and after heat treatment. EBSD was used to evaluate grain morphology, local misorientation, grain boundary characteristics, and deformation/recovery structure. Tensile testing was then correlated with the microstructural observations. The objective of this work is to clarify how post-rolling heat treatment changes the balance between recovery, grain coarsening, grain-boundary evolution, and tensile response in a multi-component Fe–15Cr ODS steel.

2. Materials and Methods

2.1 Processing conditions

The investigated alloy was selected from a machine-learning-assisted screening workflow developed for Fe-based oxide dispersion-strengthened steels. The workflow used literature-reported alloy compositions, processing parameters, and mechanical properties to identify promising composition–process combinations, with gradient-boosting-based models and SHAP analysis used to support alloy ranking [10-12]. In the present work, this workflow was used only for alloy selection; the main focus is the experimental correlation between processing, microstructure, and tensile behavior.

The selected alloy had a nominal composition of Fe–15.2Cr–6.6Mo–2.4W–1.0Zr–0.9Ti–0.35Y–0.04C, in wt.%. The alloy was designed to combine Mo/W solid-solution strengthening with oxide-forming additions based on Y, Ti, and Zr, which are commonly used to promote oxide-rich particle formation in Fe-based ODS steels [13-15]. Two material conditions were examined: the hot-rolled condition, obtained after hot isostatic pressing (HIP) consolidation and hot rolling, and the heat-treated condition, obtained after post-rolling heat treatment at 1100 °C for 24 h in Ar followed by air cooling.

2.2 Powder processing and thermomechanical treatment

A Fe–15Cr pre-alloyed powder with a particle-size range of approximately 40–130 μm was used as the base powder. Elemental powders were added to adjust the alloy composition to the nominal chemistry described above. Y₂O₃ powder with an average particle size of approximately 10 nm was introduced as the oxide source for ODS particle formation.

Mechanical alloying was carried out in stainless-steel vials using stainless-steel balls under an Ar atmosphere. The ball-to-powder weight ratio was fixed at 10:1. The milling schedule consisted of a cyclic program with a total programmed duration of 80 h, including 40 h of effective high-energy milling at 550 rpm and 40 h of low-speed intervals at 240 rpm. The low-speed intervals were used to limit excessive temperature rise and reduce cold welding during mechanical alloying.

The mechanically alloyed powders were consolidated by hot isostatic pressing at 1150 °C under 100 MPa for 3 h, followed by air cooling. The consolidated billet was then hot rolled at 1200 °C for 7 passes, resulting in a total thickness reduction of approximately 70%. After hot rolling, part of the material was retained as the hot-rolled condition, while the remaining material was heat treated at 1100 °C for 24 h under an Ar atmosphere, followed by air cooling. The detail index is presented in Table 1.

2.3 EBSD characterization

Electron backscatter diffraction was used to characterize the grain structure, local orientation gradients, and grain-boundary characteristics of the hot-rolled and heat-treated conditions. EBSD specimens were sectioned from the longitudinal rolling direction–normal direction (RD–ND) plane, with the observation direction parallel to the transverse direction. This section was selected to examine the deformation structure developed during hot rolling.

The specimens were mechanically ground and polished, followed by final surface finishing to reduce preparation-induced deformation. EBSD data were acquired using a field-emission scanning electron microscope equipped with an EBSD detector. The step size was 30 nm for the hot-rolled condition and 0.3 μm for the heat-treated condition.

The EBSD analysis included inverse pole figure (IPF) maps, grain-size maps, kernel average misorientation (KAM) maps, grain orientation spread (GOS), grain reference orientation deviation (GROD), and grain-boundary misorientation distributions. Low-angle grain boundaries (LAGB) and high-angle grain boundaries (HAGB) were classified using misorientation ranges of 2–15° and >15°, respectively [16]. Because the two EBSD datasets were acquired with different step sizes, local misorientation parameters such as kernel average misorientation and LAGB fraction were used mainly as qualitative indicators of local lattice distortion and stored deformation. The heat-treatment response was therefore interpreted from the combined trends in inverse pole figure (IPF) morphology, grain-boundary character, KAM, GOS, and GROD, rather than from direct comparison of absolute KAM or LAGB values alone.

2.4 Tensile testing

Tensile testing was performed to evaluate the mechanical response of the investigated alloy. The tensile axis was aligned parallel to the rolling direction. Sub-size sheet-type specimens followed the ASTM E8/E8M geometry, with a gauge length of 25 mm, a gauge width of approximately 5.99 mm, and a thickness of approximately 1.49 mm [17]. The initial grip separation was 80 mm, and a contact extensometer with a 25 mm gauge length was used to record strain during testing. Room-temperature tensile tests were performed for both the hot-rolled and heat-treated conditions, whereas elevated-temperature tensile testing at 650 °C was limited to the hot-rolled condition. For the 650 °C test, the specimen was heated to the target temperature and held for 20 min before loading to stabilize the gauge-section temperature [18]. For each testing condition, multiple specimens were tested where material availability allowed, typically three specimens. The 0.2% offset yield strength, ultimate tensile strength, and total elongation were extracted from the tensile test data. The tensile results were then correlated with EBSD observations to evaluate the effect of post-rolling heat treatment on the room-temperature strength–ductility balance and the elevated-temperature response of the hot-rolled condition.

3. Results

The results are organized around the processing sequence. The hot-rolled condition is first described as the baseline state before heat treatment. The microstructural changes after annealing at 1100 °C for 24 h are then examined, followed by a focused analysis of local misorientation, grain-boundary character, and tensile behavior at room temperature and 650 °C.

3.1 Microstructure of the hot-rolled condition before heat treatment

The hot-rolled condition was first examined by EBSD to establish the baseline microstructure before heat treatment. The corresponding IPF map, phase map, and grain-size distribution are shown in Fig. 1. The IPF map in Fig. 1a shows the baseline microstructure of the selected ODS alloy in the hot-rolled condition before heat treatment. Most grains were below 1 µm, but the grain morphology was not uniform across the mapped area. Very fine grains coexisted with locally coarser regions, indicating a heterogeneous hot-rolled microstructure. A strong global grain elongation along the rolling direction was not evident at this EBSD scan scale. Instead, the structure appeared fragmented, with local clusters of refined grains rather than a continuous banded morphology. As shown in the phase map in Fig. 1b, the sample consisted mainly of a body-centered cubic (BCC) ferritic matrix, with a BCC-indexed fraction of 0.962. A minor face-centered cubic (FCC)-indexed fraction of 0.038 was also detected. These minor regions are described only as FCC-indexed regions, because EBSD phase indexing alone cannot determine their chemistry or physical identity. No specific assignment to oxide, carbide, or retained austenite is made at this stage.

The grain-size histogram in Fig. 1c confirms the refined nature of the hot-rolled condition. The number-average grain diameter was approximately 0.13 µm, while the area-average grain diameter was approximately 0.45 µm. The histogram is strongly weighted toward the smallest grain-size range, showing that ultrafine grains dominated the number fraction. At the same time, the larger area-average value indicates that a smaller population of coarser grains occupied a noticeable fraction of the mapped area. These values are compared with the heat-treated condition in Table 2.

3.2 Microstructural evolution after heat treatment

Heat treatment at 1100 °C for 24 h strongly altered the grain structure established during hot rolling. In the heat-treated condition, the dense ultrafine microstructure described in Section 3.1 was replaced by a much coarser matrix. As shown in the IPF map in Fig. 2a, the matrix consists mainly of larger equiaxed or semi-equiaxed grains, rather than the fragmented ultrafine grain clusters observed in the hot-rolled condition. The extent of grain growth is summarized in Table 2. The number-average grain diameter increased from approximately 0.13 µm in the hot-rolled condition to approximately 3.22 µm after heat treatment. The area-average grain diameter also increased sharply, from approximately 0.45 µm to approximately 11.25 µm. This corresponds to an approximately 25-fold increase in both grain-size measures. Although the hot-rolled and heat-treated EBSD datasets were acquired with different step sizes, the magnitude of this increase is large enough to confirm substantial grain coarsening during the 1100 °C/24 h treatment.

The large gap between the number-average and area-average grain sizes persisted after heat treatment. This divergence indicates that grain coarsening was not fully uniform. Instead, a limited population of coarse grains occupied a large fraction of the mapped area, while smaller grains still remained. This behavior suggests selective grain coarsening and may be related to abnormal grain-growth behavior reported in nanostructured ferritic steels during high-temperature treatment [19]. However, this interpretation is based mainly on grain-size statistics. The role of boundary character is examined later through EBSD misorientation analysis. The phase map in Fig. 2b gives an almost fully BCC-indexed matrix after heat treatment. The BCC-indexed fraction increased to 0.998, while the minor FCC-indexed fraction decreased from 0.038 to 0.002, as listed in Table 2. This decrease should be treated as an EBSD indexing result, not as direct proof of phase dissolution or transformation. The difference in EBSD step size is important here: the hot-rolled condition was acquired with a 30 nm step size, whereas the heat-treated condition was acquired with a 0.3 µm step size. Since EBSD phase detection depends on spatial resolution, nanoscale or sparsely distributed FCC-indexed regions may be under-sampled in the heat-treated map acquired with the larger step size [20]. Thus, the main microstructural change after heat treatment was the transition from an ultrafine, heterogeneous hot-rolled structure to a coarser and BCC-dominated heat-treated structure. The grain-growth behavior appears selective rather than fully uniform. Whether this coarsened state also involves recovery, substructure relaxation, and changes in grain-boundary character is addressed in Section 3.3.

3.3 Orientation-spread and grain-boundary evolution

The grain coarsening described in Section 3.2 was accompanied by a clear relaxation of internal orientation gradients. Kernel average misorientation (KAM), grain reference orientation deviation (GROD), grain orientation spread (GOS), and grain-boundary misorientation were examined together to determine whether the heat-treated condition still retained the hot-rolling substructure. Here, GOS was used as a grain-scale descriptor of intragranular orientation spread, whereas LAGB/HAGB fractions were used as boundary-scale descriptors of the grain-boundary network. Because local misorientation values are sensitive to EBSD acquisition conditions, especially step size, the analysis focuses on trends shared by several EBSD descriptors rather than on absolute KAM or LAGB values alone [21, 22].

In the hot-rolled condition, the KAM and GROD maps in Figs. 3a and 3c show that the microstructure was not strain-free after rolling. High-KAM regions appeared locally near grain boundaries, triple junctions, and selected intragranular areas. The maximum GROD value reached approximately 16.8°, indicating clear intragranular orientation gradients. These features are consistent with retained deformation substructure after hot rolling. The GOS map and partition in Fig. 3e describe this substructure at the grain scale. In the hot-rolled condition, the GOS partition showed 64.1% in the 0–2° range, 33.5% in the 2–5° range, and 2.2% above 5° [23-25]. This distribution suggests that part of the hot-rolled matrix had already reached a relatively low orientation-spread state, while a substantial fraction still retained internal orientation spread. Therefore, the hot-rolled condition should be described as a recovered/substructured state rather than a fully recrystallized structure. GOS-based partitioning can support recrystallization analysis, but the selected threshold should be interpreted with the alloy system and processing history in mind [25].

The grain-boundary statistics in Table 3 describe the boundary network separately from the GOS partitioning. In the hot-rolled condition, the relative LAGB fraction was 0.16, whereas the relative HAGB fraction was nearly 0.46. The boundary misorientation analysis also showed a pronounced low-angle peak near 2.5°, with a fraction of approximately 0.49, and the average boundary misorientation angle was about 21.6°. This low-angle peak indicates that a large part of the classified boundary network was still associated with deformation-induced subgrain boundaries.

After heat treatment, the internal orientation spread decreased markedly. The heat-treated condition showed lower overall KAM contrast in Fig. 3b and weaker intragranular orientation gradients in Fig. 3d. The maximum GROD value decreased from approximately 16.8° to approximately 6.9°. The GOS partition also shifted strongly toward the low-spread range: the 0–2° partition increased from 0.64 to 0.99, the 2–5° partition decreased from 0.34 to 0.01, and the >5° partition became nearly absent. These changes are summarized in Table 3 and are consistent with substantial relaxation of the hot-rolling deformation substructure during heat treatment.

The boundary network also changed after heat treatment. The relative LAGB fraction decreased from 0.16 to 0.06, while the relative HAGB fraction increased from 0.46 to 0.87. At the same time, the low-angle peak near 2.5° decreased from approximately 0.49 to 0.08, and the average boundary misorientation increased from about 21.6° to 37.1°. This shift indicates that the heat-treated condition became HAGB-dominated and contained much less low-angle substructure than the hot-rolled condition. The result is consistent with substructure reduction, recovery, and recrystallization-like boundary evolution during heat treatment, but it does not prove complete recrystallization.

Collectively, the EBSD results show that heat treatment did more than coarsen the grains. It reduced intragranular orientation spread, weakened the low-angle deformation substructure, and shifted the boundary network toward high-angle boundaries. The heat-treated alloy is therefore best described as a coarser, strongly recovered, HAGB-dominated state. The following section relates this EBSD-scale microstructural evolution to the tensile response.

3.4 Tensile properties

The tensile response provides the mechanical counterpart to the microstructural evolution described above. As shown in Table 4, the hot-rolled condition exhibited very high room-temperature strength but limited ductility. The yield strength reached 820 MPa, while the ultimate tensile strength increased to 2079 MPa, with a total elongation of 2.3%. The apparent increase from yield strength to ultimate tensile strength was approximately 1259 MPa. Because the total elongation was limited, this increase should not be interpreted as a conventional strain-hardening capacity alone. The high room-temperature strength is more reasonably associated with the combined effects of ultrafine grains, retained deformation substructure, alloying-related matrix strengthening discussed in Sections 3.1–3.3. The individual contribution of each factor cannot be separated from the present tensile data. At 650 °C, the hot-rolled condition showed lower strength but much higher ductility. The yield strength decreased to 451 MPa and the ultimate tensile strength decreased to 571 MPa, while the total elongation increased to 34.9%. This response indicates strong thermal softening during elevated-temperature deformation. The large increase in elongation suggests that plastic strain was accommodated more readily at 650 °C than at room temperature. This behavior may reflect thermally activated plasticity and dynamic recovery or restoration processes promoted by the high stored energy of the hot-rolled substructure. Post-rolling heat treatment changed the room-temperature strength–ductility balance. After heat treatment at 1100 °C for 24 h, the yield strength decreased from 820 MPa to 680 MPa, and the ultimate tensile strength decreased from 2079 MPa to 935 MPa. In contrast, the total elongation increased from 2.3% to 13.9%. This trend agrees with the EBSD results in Sections 3.2 and 3.3, where heat treatment caused substantial grain coarsening, reduced local orientation gradients, weakened the low-angle boundary contribution, and produced a more recovered, HAGB-dominated structure.

These tensile results show that heat treatment improved room-temperature ductility at the cost of strength. The hot-rolled condition provided the highest strength but fractured with limited plastic strain, whereas the heat-treated condition showed lower strength and improved elongation. The 650 °C test further shows that the hot-rolled alloy can accommodate substantially more plastic strain at elevated temperature, although elevated-temperature tensile data for the heat-treated condition are not included in the present dataset.

4. Discussion

4.1 Effect of heat treatment on deformation–recovery structure

Heat treatment changed the investigated alloy in two coupled ways: it coarsened the ferritic grain structure and relaxed much of the deformation substructure retained after hot rolling. The hot-rolled condition was already partially recovered rather than simply as-deformed, as indicated by the fragmented grain morphology, the coexistence of low-angle and high-angle boundaries, and the presence of a low-GOS grain population. Even so, the persistence of local orientation gradients and a distinct low-angle boundary component shows that a substantial stored-deformation substructure remained after rolling. This intermediate state is important because it defines the starting point for the subsequent annealing response.

The EBSD descriptors are most reliable when read together. KAM-based and related orientation-spread metrics are useful for assessing stored plastic deformation, but their absolute values depend on step size, kernel definition, measurement noise, and grain size [21, 22]. For that reason, the present interpretation is based on the consistent trend among KAM, GROD, GOS, and boundary statistics rather than on any single parameter. After heat treatment at 1100 °C for 24 h, the decrease in local misorientation, the lower LAGB fraction, the shift of the GOS distribution toward lower spread, and the higher HAGB fraction all point to a coarser, strongly recovered, HAGB-dominated structure. In this context, the heat-treated condition is better described as strongly recovered than as a fully discontinuously recrystallized state [25].

This structural relaxation provides the most direct explanation for the room-temperature strength–ductility trade-off shown in Table 4. In the hot-rolled condition, grain-boundary strengthening and deformation-substructure strengthening were active at the same time. After heat treatment, both contributions were reduced: grain growth weakened the boundary contribution, while recovery removed part of the internal defect structure retained after rolling. By contrast, the lower internal orientation gradients and reduced substructure density made plastic strain easier to accommodate. The heat-treated condition was therefore softer, but mechanically more compliant at room temperature.

4.2 Relationship between microstructure and tensile behavior

The tensile response in Table 4 can be read as the mechanical expression of this microstructural balance. In the hot-rolled condition, high strength likely arose from the combined action of ultrafine grains, retained deformation substructure, and alloying-related matrix strengthening. At the same time, the microstructure was not uniform. The difference between number-average and area-average grain size, together with the retained intragranular orientation gradients, indicates a heterogeneous deformation-resistant matrix rather than a fully equilibrated ultrafine-grained structure. That heterogeneity can promote strain partitioning and local stress concentration during room-temperature loading, which helps explain why very high strength coexisted with limited tensile ductility. This point is important because the high tensile resistance of the hot-rolled condition should not be assigned to grain refinement alone. A grain-size effect is clearly present, but it operated alongside retained deformation substructure and matrix-strengthening effects from the multi-component alloy design. The current results support that multi-factor picture, even though the individual contributions cannot be separated quantitatively from the available data.

Heat treatment shifted the balance in the opposite direction. As the ferritic matrix became coarser and more strongly recovered, the alloy lost part of the microstructural barriers that had sustained high strength in the hot-rolled state. In return, the reduced internal gradients and lower low-angle boundary density allowed more stable strain accommodation at room temperature. The result was the strength–ductility trade-off expected from this microstructural evolution: lower strength, but improved elongation.

4.3 Implications and remaining limitations

The primary implication of these findings is that the property balance of the investigated alloy is highly sensitive to post-rolling thermal exposure, indicating a narrow processing window. The alloy-screening strategy was effective in identifying a high-strength Fe-based ODS candidate, yet the final property balance was governed by more than composition alone. In ODS systems, the final property balance is generally governed by the combined effects of alloy chemistry, grain/subgrain structure, and thermomechanical history [2–5]. The present results fit that framework well. In practical terms, hot rolling generated a high-strength condition by retaining a fine, substructured ferritic matrix. The selected post-rolling heat treatment improved room-temperature ductility, but it also erased a large part of the microstructural refinement responsible for that strength. The investigated alloy should therefore be regarded as a partially stabilized ODS candidate rather than an optimized thermally stable microstructure.

The main remaining limitation is that the present dataset does not directly resolve the nanoscale oxide population or its evolution during heat treatment. Therefore, the discussion of microstructural stability is limited mainly to EBSD-scale grain/subgrain evolution and tensile response. Future work should examine particle-scale stability and elevated-temperature tensile behavior of the heat-treated condition to clarify the full processing–microstructure–property relationship.

5. Conclusion

This study examined the microstructural evolution and tensile response of a machine-learning-selected, multi-component Fe–15Cr-based ODS steel after hot rolling and post-rolling heat treatment. The main conclusions are as follows:

1. The hot-rolled condition exhibited a BCC-dominated, ultrafine, and heterogeneous ferritic microstructure. EBSD analysis showed a number-average grain size of approximately 0.13 μm and an area-average grain size of approximately 0.45 μm, indicating that fine grains dominated by number while locally coarser regions occupied a noticeable mapped area.

2. Heat treatment at 1100 °C for 24 h caused substantial grain coarsening. The number-average and area-average grain sizes increased to approximately 3.22 μm and 11.25 μm, respectively, corresponding to an approximately 25-fold increase. The heat-treated condition also became more strongly BCC-indexed, although the decrease in minor FCC-indexed regions should be interpreted cautiously because of the different EBSD step sizes used for the two conditions.

3. Local misorientation and grain-boundary analyses showed that heat treatment relaxed much of the hot-rolling deformation substructure. The heat-treated condition showed reduced KAM/GROD contrast, a lower LAGB fraction, a strong shift of the GOS distribution toward low orientation spread, and a higher HAGB fraction. These results indicate a transition from a recovered but still substructured hot-rolled matrix to a coarser, strongly recovered, HAGB-dominated state.

4. The tensile response reflected the competition between strengthening and strain accommodation. The hot-rolled condition showed very high room-temperature strength, with a yield strength of 820 MPa and an ultimate tensile strength of 2079 MPa, but limited elongation of 2.3%. At 650 °C, the hot-rolled alloy showed lower strength but much higher elongation of 34.9%. After heat treatment, room-temperature strength decreased, while elongation increased to 13.9%. This confirms that post-rolling heat treatment improved ductility at the cost of strength.

Article information

Funding

This research was supported by the NANO & Material Technology Development Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT (No. RS-2024–00402289). This work was also supported by the Korea Planning and Evaluation Institute of Industrial Technology (KEIT) grant funded by the Ministry of Trade, Industry & Energy (MOTIE) of the Republic of Korea (No. 2410016473).

Conflict of Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. In addition, Prof.J. Kim serves as an Editor of the Journal of Powder Materials but had no involvement in the peer-review or decision-making process for this manuscript.

Data Availability Statement

Data will be made available on request.

Author Information and Contribution

Trung Thanh Pham: Writing – original draft, Visualization, Methodology. Woo-Hyeok Kim: Investigation, Data curation. HyunCheol Kim: Validation, Jeoung Han Kim: Supervision, Project administration

Acknowledgments

None.

Fig. 1.

EBSD microstructure of the selected candidate alloy in the hot-rolled condition (a) IPF map, (b) EBSD phase map and (c) grain-size histogram.

Fig. 2.

EBSD microstructure of the selected candidate alloy after post-rolling heat treatment: (a) IPF map, (b) EBSD phase map, and (c) grain-size histogram with summary statistics.

Fig. 3.

Local misorientation and orientation-spread maps of the selected ODS alloy: (a, b) kernel average misorientation (KAM) maps, (c, d) grain reference orientation deviation (GROD) maps, and (e, f) grain orientation spread (GOS) maps. Panels (a, c, e) correspond to the hotrolled condition, and panels (b, d, f) correspond to the heat-treated condition.

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Figure & Data

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