Microstructure and properties of oxide dispersion-strengthened alloys

Room temperature mechanical properties

The strengthening of ODS steels primarily arises from grain refinement, oxide particle dispersion, and dislocation interactions. Reducing oxide particle size and increasing their number density are common methods to improve the properties of ODS steels. Nano-oxides act as pinning sites for dislocations, enhancing strength while maintaining adequate ductility. The grain size, dislocation density, and carbide precipitation further influence mechanical properties. Alloying elements such as Cr, Al, Mn, Ti, and Zr modify the microstructure and influence mechanical responses. Typically, the introduction of nano-oxide particles can increase the hardness, yield strength (YS), and ultimate tensile strength (UTS) of the material, but it can also reduce the material's ductility.

Effects of alloying elements on room-temperature mechanical properties of ODS alloys

Alloying additions in ODS alloys primarily influence room-temperature mechanical properties through their effects on oxide nanoparticle chemistry, size, number density, and the resulting phase balance within the matrix. While many nanoparticle-forming elements contribute to strengthening, their macroscopic effects vary significantly with elemental type, concentration, and interactions with the base alloy system. Zhou et al. ⁴⁷ investigated 9Cr-based ODS martensitic steels strengthened with either Zr or Hf and processed by mechanical milling, discharge plasma sintering, and subsequent tempering at 800 °C. The Zr-containing alloy exhibited an average microhardness, ultimate tensile strength (UTS), and yield strength (YS) of 316 HV, 1098 MPa, and 918 MPa, respectively. In contrast, the Hf-containing alloy reached 413 HV, 1203 MPa, and 963 MPa. At low addition levels (0.1 wt.%), Hf was more effective in increasing hardness and strength. Microstructural analysis indicated that both Zr and Hf additions reduced the ferrite fraction compared with Al- and Ti-containing ODS steels.^74,75 In the Zr- and Hf-modified alloys, ferrite was identified as residual ferrite without M₂₃C₆ precipitation or as transformed ferrite, respectively. The 9Cr–Hf ODS steel exhibited higher strength but reduced ductility relative to the Zr-containing counterpart, a trade-off attributed to finer nanoparticle size and a reduced density of ferrite–martensite interfaces, which limited microstructural heterogeneity and crack initiation sites.

Peng et al. ⁷⁶ studied the effect of multiple principal elements on the mechanical properties of ODS alloys. They experimentally prepared equiatomic ternary CrFeNi, quaternary CoCrFeNi, and quintary CoCrFeNiMn alloys with 1 wt.% Y₂O₃ and 1 wt.% Zr by MA and SPS. An increase in the number of principal elements resulted in a reduction in the hardness and compressive YS. The hardness, YS, and fracture strain of the ternary alloy were 475 HV, 1531 MPa, and 20%, respectively. The five-component CoCrFiNiMn ODS alloy has the lowest hardness and compressive YS of 1228 MPa. Grain-boundary strengthening plays a major role in this process, and microscopic analysis shows that, as the number of principal elements increases, grain size and the average size of nano-oxides increase. In contrast, the number and density of oxides decrease. When the addition of reinforcing nanoparticle-forming elements exceeds a certain threshold, it affects the material's mechanical properties at room temperature. An Fe-24Mn-0.4Ti-0.25Y₂O₃ (24Mn) alloy prepared by MA and field-assisted sintering by Seils et al. ⁷⁰ has a higher YS than Fe-34Mn-0.4Ti-0.25Y₂O₃ (34Mn) with higher Mn content, which are 1081 ± 2 MPa and 1055 ± 5 MPa, respectively. Unlike other additives, although Al can also form Y-Al-O complexes, its larger particle size does not improve the material's UTS or YS. Wang et al. ⁵⁹ compared 9Cr-ODS steel with a counterpart containing 3 wt—% Al. The UTS and YS of the alloy decreased from 968 ± 8 MPa and 836 ± 7 MPa to 942 ± 13 MPa and 791 ± 7 MPa, respectively, but the elongation increased from 18 ± 0.8% to 23.5 ± 1.2%. Table 1 shows the mechanical properties of some of the ODS steels from the literature. Yttrium is one of the most important and commonly used elements for forming nano-oxides in ODS alloys. Other elements, such as Al and Ti, can react with Y₂O₃ to form nanoxides. These oxide particles can effectively impede dislocation motion and enhance creep strength. The addition of Al can improve oxidation resistance; however, excessive amounts may reduce oxide dispersion uniformity and affect strengthening.

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Table 1.

ODS steel mechanical properties in the literature.^45,77–90

Materials	Hardness [HV]	Tensile strength [MPa]	Yield strength [MPa]	Elongation [%]
Fe-5Cr-4.5Al-0.6Zr 45	–	875	–	21
14Cr-2W-0.3Ti-0.3Y2O3 78	448	1420	1340	–
Conventional 316L(N) 87	–	601	306	–
Ni46HESA 1 vol%Y2O3 88	–	1155	–	–
Fe–18Cr–8Ni–3Cu-0.4Nb-0.35Y2O3 89	–	995	700	–
Fe-9Cr-1.5W-0.2Ta-0.2V-0.4Mn-0.3Si-0.3Y2O3 90	–	1459	1315	–
Fe-17.7Cr-3.84Al-2.13W-0.59Ti-0.05Ta-0.19Y-0.17O 79	–	-	–	13
18Cr ferritic steel 80	396 – Axial 402– ransverse	–	–	-
Fe-14Cr-1.5W-0.9Zr-0.45Y2O3 77	–	994.3	–	13
Fe-22Cr-3Al 82	–	797	669	9
Fe-9Cr-1.1W-0.4Mn-0.2V-0.12Ta-0.3Y2O3 （wt.%） 83	–	734	–	–
Eurofer 97 84		680	–	–
Fe-22Cr-5A1-3Mo-0.19 (Ti + Y + Zr) 0.01O/0.03O 85	–	–	498 / 603	–
ODS 316L 86	–	723	458	–

Ti can react with Y₂O₃ to form Y-Ti-O oxide nanoparticles, such as Y₂Ti₂O₇, YTiO₃, etc., which are usually smaller in size than Al-containing oxide particles, more evenly distributed, and have a better dispersion strengthening effect. Cr can improve oxidation and corrosion resistance, making the alloy more suitable for high-temperature environments. Still, an excessive amount may cause brittle phases (e.g., Cr₂O₃) to precipitate, thereby reducing ductility. Similarly, the addition of appropriate amounts of W and Mo can improve the strength of ODS alloys, but excessive concentrations can lead to brittleness. The addition of Zr, Hf, V, and other elements has also been shown to improve the radiation resistance of materials. In addition, the appropriate addition of Re and C can improve the material's ductility; however, when C is added in excess, it forms carbides such as M₃C and M₂₃C₆, leading to grain-boundary embrittlement.

Additionally, the fabrication method of ODS steels affects the materials’ mechanical properties at room temperature to some extent. For example, hot rolling can make the material's microstructure anisotropic. Oksiuta et al. ⁹¹ found that 14Cr ODS steel exhibits a higher dislocation density and a finer microstructure in the rolling direction, which has no evident effect on hardness changes in the longitudinal and vertical directions but significantly affects the material's elongation in different directions. ⁹² Evenly dispersed nano-scale oxide particles can be obtained by SPS. Samples prepared by SPS typically exhibit high strength and toughness; however, their dimensions and shape are limited. ODS steels manufactured by powder hot extrusion usually have finer grains and higher strength, but poor plasticity. HIP can fabricate parts with complex shapes and large dimensions, and the resulting materials exhibit good strength and ductility.

Elevated temperature mechanical properties

Mechanical properties such as strength, ductility, creep resistance, and toughness are essential requirements that should be tested and assessed before selecting a material for a particular application. The material selected must retain good mechanical properties at the elevated temperatures required for nuclear applications.

Creep behavior

Dymácek et al. ¹⁰⁶ observed that the FeAlOY ODS alloys possess exceptionally high-temperature creep resistance. They concluded that these alloys exhibited exceptional creep resistance up to 1300 °C. The enhanced performance of these FeAlOY ODS alloys is due to higher Y content, which enables secondary recrystallization. Ohtsuka et al. ¹⁰⁷ studied the creep-rupture response of 9Cr ODS steel containing different mean particle distances between the oxide particles at 700 °C. It was observed that specimens with higher mean particle distances exhibit higher secondary creep strain rates and lower creep-rupture strengths. Hence, dispersion strengthening decreases as the mean particle distance increases. Interestingly, in ODS alloys, creep strength was also proportional to Vickers hardness. Yano et al. ¹⁰⁸ reported that 9Cr ODS steel with a tempered martensitic microstructure performs better under creep loading at 1000 °C than other alloys, such as 11Cr-ferritic/martensitic steels. Creep test of 18Cr ODS steel ¹⁰⁹ at 650 °C also revealed lower creep strain and limited tertiary creep. TEM analysis showed that the ability of nanoparticles to pin dislocations at higher temperatures is compromised by thermal activation. The authors highlighted that conducting in situ deformation experiments could reveal greater details of the deformation behavior in these alloys. Unlike ferritic ODS steels, ¹⁰⁰ the tertiary creep dominates in austenitic SS316L ⁸⁶ ODS steel and consumes more than 50% of the total life, as shown in Figure 5. However, the creep-rupture life of the austenitic SS316L ODS steel was much shorter, with a lower stress exponent, indicating limited interaction between dislocations and nanoparticles.

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Figure 5.

Creep strain at a temperature of 650 °C in austenitic and ferritic ODS alloys at stress levels of 200 and 250 MPa, respectively. Reproduced from Ref. SS316L, ⁸⁶ and 1400 h. ¹⁰⁹

The creep behavior of 304L ODS steel fabricated via LPBF was evaluated at 700 °C. The lifetime to rupture under a 70 MPa load was 1400 h. ¹⁰⁹ However, at a stress of 100 MPa, the specimen lasted only for 153 h. A stress exponent of 7.7 was obtained, indicating that dislocation climb is the dominant creep mechanism. The 304L ODS steel exhibited creep rates two orders of magnitude lower under the same stress conditions compared to conventional 304L steel. In this alloy, the creep life is also dominated by the tertiary creep.

As a result, ODS alloys exhibit excellent high-temperature creep resistance, with ferritic variants minimizing tertiary creep, while austenitic ODS alloys show shorter rupture life and dominant tertiary creep. The excellent creep resistance of ODS alloys is primarily attributed to the role of oxide dispersions in reducing dislocation creep and diffusional creep. Factors such as Y content, particle spacing, and alloying elements also influence creep behavior, while thermal activation limits dislocation pinning at extreme temperatures.

Beyond dispersoid characteristics, the evolution of grain structure during processing and high-temperature exposure also plays a critical role in determining creep performance. Secondary recrystallization is a key microstructural factor governing the high-temperature creep behavior of mechanically alloyed ODS alloys. While ultrafine-grained structures can provide high strength, they are generally detrimental under creep conditions due to enhanced grain-boundary sliding and diffusional creep. In contrast, coarse-grained microstructures formed through secondary recrystallization significantly reduce grain-boundary area and improve creep resistance.

For example, Bártková et al. demonstrated that an Fe–Al–O ultra-fine-grained nanocomposite with an initial grain size of ∼150–200 nm undergoes abrupt grain coarsening to ∼50 μm after annealing at 1000 °C for 8 h, driven by the rapid growth of a limited number of dislocation-free grains. ¹¹⁰ More recent work on Fe–Al–Cr-based ODS alloys further shows that grain size can evolve from ∼100 nm after hot consolidation to 1–10 mm following annealing, with abnormal grain growth occurring rapidly after an incubation period that is highly sensitive to processing parameters such as deformation and yttrium content. ¹¹¹ In addition, Svoboda et al. reported that oxide chemistry influences both secondary recrystallization and oxide coarsening behavior in Fe–10Al–4Cr-based ODS alloys, highlighting the importance of dispersoid composition in stabilizing microstructures at elevated temperatures. ¹¹²

These studies demonstrate that creep resistance depends not only on dispersoid stability and spacing, but also on achieving a fully recrystallized coarse-grained microstructure rather than a mixed ultrafine/coarse-grained structure, which can accelerate creep deformation. The competition between stored dislocation energy, capillary driving forces, oxide coarsening, and Zener pinning ultimately controls secondary recrystallization and, consequently, the long-term creep performance of ODS alloys.

Fracture toughness and failure mechanisms

Fracture toughness in ODS alloys is governed by a complex interplay between dispersoid characteristics, grain structure, alloy composition, and processing-induced anisotropy. In contrast to conventional steels, the presence of a high number density of nanoscale oxide particles and elongated grain morphologies introduces additional mechanisms that influence crack initiation and propagation, particularly across a wide temperature range.

The ductile-to-brittle transition temperature (DBTT) in ODS alloys is highly sensitive to both dispersoid population and matrix composition. Oksiuta et al. reported DBTT values of 77 °C and −24 °C for Y₂O₃- and Fe₂Y-strengthened ODS steels, respectively, indicating that reduced nanoparticle volume fraction and modified particle chemistry can improve low-temperature toughness. ¹¹³ Similarly, increases in Cr, W, and Ti content have been shown to raise DBTT and reduce low-temperature toughness, reflecting the combined influence of solid-solution strengthening and matrix hardening on crack-tip plasticity. ¹¹⁴ In addition, higher Y₂O₃ contents, while beneficial for high-temperature strength, may reduce impact toughness, highlighting the need to balance dispersoid strengthening with fracture resistance. These effects directly influence both crack initiation and early-stage crack propagation by controlling the extent of plastic deformation at the crack tip and the susceptibility to localized damage accumulation.

A key characteristic of ODS alloys is their anisotropic fracture behavior, which arises from elongated grain structures and crystallographic texture introduced during thermomechanical processing. Das et al. ¹¹⁵ demonstrated that fracture toughness in hot-rolled 13Cr ODS steel depends strongly on crack orientation, with higher toughness observed in the L–T orientation compared to T–L. This behavior is attributed to enhanced crack branching and secondary cracking when the crack propagates perpendicular to elongated grains, whereas crack propagation parallel to grain elongation is facilitated by reduced resistance along aligned microstructural features. These observations highlight the critical role of grain morphology and texture in controlling crack path stability and energy absorption mechanisms. Such anisotropic crack propagation characteristics are also closely related to fatigue crack growth behavior, where similar microstructural features govern cyclic damage accumulation and crack advancement.

Thermomechanical processing further modifies fracture behavior by altering grain structure and phase distribution. For example, in 9Cr ODS steels, water-quenched specimens exhibit improved toughness over a wide temperature range compared to air-cooled conditions, primarily due to the presence of residual ferrite phases that promote secondary cracking and energy dissipation. ¹¹⁶ In addition, bimodal grain structures and fine, uniformly distributed oxide nanoparticles contribute to enhanced crack-tip blunting and resistance to crack propagation.

From a mechanistic perspective, fracture behavior in ODS alloys is strongly influenced by dispersoid–dislocation interactions and grain boundary characteristics. Fine, coherent nanoclusters can promote more homogeneous plastic deformation and delay crack initiation by impeding dislocation motion. In contrast, coarser or inhomogeneously distributed particles may act as local stress concentrators, facilitating void nucleation and accelerating crack propagation. Furthermore, grain boundaries in ODS alloys can act as both barriers to crack growth and preferential sites for damage accumulation, depending on their cohesion and interaction with dispersoids. In terms of failure modes, ODS alloys exhibit a transition from ductile to brittle fracture depending on temperature, microstructure, and alloy composition. At higher temperatures, fracture is typically characterized by ductile mechanisms involving microvoid nucleation, growth, and coalescence, facilitated by dislocation activity and plastic deformation around dispersoids. At lower temperatures, limited dislocation mobility promotes brittle fracture, often associated with cleavage or quasi-cleavage mechanisms. In addition, intergranular fracture may occur in cases where grain boundary cohesion is reduced, particularly in alloys with unfavorable grain boundary chemistry or coarse dispersoid distributions. The competition between these failure mechanisms is strongly influenced by dispersoid stability, grain structure, and matrix composition, and directly governs the observed ductile-to-brittle transition behavior. In summary, fracture toughness in ODS alloys is governed by coupled microstructural factors rather than a single controlling parameter, with dispersoid stability, grain morphology, alloy composition, and processing history collectively determining crack initiation, propagation pathways, and the ductile-to-brittle transition behavior. These interactions underscore the importance of microstructural design in optimizing fracture resistance for extreme environment applications.

To facilitate comparison across the different ODS alloy systems discussed above, a representative summary of key microstructural parameters and mechanical properties is provided in Table X. While the reported values vary depending on alloy composition, processing route, and testing conditions, the table highlights general trends linking dispersoid characteristics, matrix structure, and mechanical performance. In particular, it illustrates the distinct strengthening responses of ferritic, ferritic–martensitic, austenitic, and Ni-based ODS alloys, and emphasizes the critical role of nanoscale oxide dispersions in governing strength–ductility balance and high-temperature performance. Table 2 shows the comparative summary of microstructure–property relationships in ODS alloys. Unlike Table 1, which compiles reported mechanical property values, Table 2 provides a comparative summary linking explicitly reported microstructural parameters to the corresponding mechanical response across representative ODS alloy systems. The table is intended to highlight how dispersoid size, number density, grain structure, and alloy class collectively influence strength, ductility, thermal stability, and irradiation tolerance.

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Table 2.

Comparative summary of microstructure–property relationships in ODS alloys

Alloy system	Representative alloy / processing route	Dispersoid / microstructural characteristics	Mechanical response	Key microstructural / mechanistic feature	Ref.
Ferritic ODS steel	Fe–5Cr–4.5Al–2W–0.15V–0.6Zr + 0.35 wt.% Y2O3 (5CrAlZr ODS; MA + HIP + forging)	Mean particle diameter 10 nm; inter-particle spacing 80 nm; number density 1.23 × 1022 m−3; average grain size 0.9 μm	RT: UTS 875 MPa, TE 21%; 700 °C: UTS 234 MPa, TE 58%	Mainly δ-Y4Zr3O12 nanoparticles; finer and denser particles contribute to good strength–ductility balance	78
Ferritic ODS steel	Fe–14Cr–1.5W–0.9Zr–0.45Y2O3 (14Cr-Zr-ODS; MA + HIP)	Mean particle diameter 8.19 nm; inter-particle spacing 17.54 nm; number density 2.31 × 1015 m−2; volume fraction 0.535%	RT: YS 816.3 ± 30.0 MPa, UTS 994.3 ± 9.5 MPa, UE 13.0 ± 1.3%	Zr addition refines particles and increases density relative to Zr-free 14Cr-ODS; dominant fine phase identified as Y4Zr3O12	77
Ferritic ODS steel	Fe–14Cr–2W–0.3Ti–0.3Y2O3 (hot extrusion + hot rolling + HT)	Average Y–Ti–O nanocluster size 5 ± 3.7 nm; number density 1.5 × 1023 m−3	RT (transverse): YS ∼1340 MPa, UTS ∼1420 MPa; 750 °C: YS ∼325 MPa (longitudinal) and ∼305 MPa (transverse)	Fine Y–Ti–O nanoclusters plus elongated anisotropic grain structure produce very high strength	45
High-Cr ferritic ODS steel	Fe–22Cr–3Al (STARS + HIP)	Average grain size 6.3 μm; average oxide size 97 nm; oxides mainly located at grain boundaries	RT: UTS 797 MPa, YS 669 MPa, elongation 9%; 500 °C: UTS 536 MPa, YS 425 MPa, elongation 24%	Coarser grain-boundary (Y,Al)-containing oxides give moderate strength but limited room-temperature ductility	82
Ferritic–martensitic ODS steel	Fe–9Cr–1.5W–0.2Ta–0.2V–0.4Mn–0.3Si–0.3Y2O3 (Si-containing 9Cr ODS)	Average grain size ∼410 nm; nanoparticle average diameter ∼5.8–6 nm; number density ∼1.8 × 1024 m−3	RT: YS 1315 MPa, UTS 1459 MPa; 550 °C: YS ∼800 MPa	Ultrafine grain structure and high number density of Y–Si–O nanoparticles give very high strength	90
Martensitic ODS steel	9Cr–Zr ODS (tempered)	Mean oxide size after tempering 13 nm	YS 918 MPa, UTS 1098 MPa, total elongation 11%	Zr-modified oxides give good strength–ductility combination after tempering	47
Martensitic ODS steel	9Cr–Hf ODS (tempered)	Mean oxide size after tempering 7.5 nm (nearly unchanged from 7.2 nm before tempering)	YS 963 MPa, UTS 1203 MPa, total elongation 7%	Hf gives finer and more thermally stable oxide population than Zr, with higher strength but lower ductility	47
Austenitic ODS steel	ODS 316L (MA + HIP + forging / rolling)	Average grain size 26 μm; oxide size distribution mainly 10–50 nm and 50–200 nm	RT: YS 458 MPa, UTS 723 MPa	Higher room-temperature strength than conventional 316L; particles <15 nm were coherent or semi-coherent, larger particles incoherent	86
Austenitic ODS steel	Fe–18Cr–8Ni–3Cu–0.4Nb–0.35Y2O3 (MA + HIP)	Average grain size ∼490 nm; average dispersoid size 8.77 nm	RT: UTS 995 MPa, Rp0.2 700 MPa; 650 °C: strengths around 2× those of Super304H	Ultrafine grains plus numerous Y2O3-based particles strongly improve strength; many particles show Y-rich shell / Cu-rich core association	89
Austenitic ODS steel (LPBF)	304L ODS	Nanoparticle size 25 ± 10 nm as-built; 36 ± 13 nm after 1000 °C/100 h; 38 ± 13 nm after 1100 °C/100 h; 50 ± 15 nm after 1200 °C/100 h; average grain size 7.7 ± 6.7 μm as-built and 9.1 ± 7.2 μm after 1200 °C/100 h	RT: YS 575 ± 8 MPa, UTS 700 ± 13 MPa, elongation 32 ± 5%; 600 °C: YS 290 ± 2 MPa, UTS 370 ± 4 MPa; 800 °C: YS 145 ± 5 MPa, UTS 152 ± 6 MPa	Retained cellular substructure and thermally stable Y–Si–O nanoparticles hinder recrystallization and grain growth	104
Ni-rich ODS alloy	Ni46 HESA + 1 vol.% Y2O3 (PM 1Y)	Average grain size 525 nm; γ′ volume fraction 39.6%; average γ′ size 265 nm; Y2O3 volume fraction 0.93%; Y2O3 size 23 nm	RT: YS 1155 MPa; 700 °C: YS 880 MPa; 800 °C: YS 485 MPa	Grain-boundary plus oxide-dispersion strengthening markedly improves high-temperature strength	88
Ni-based ODS alloy under irradiation	PM1000-type Ni-based ODS	Unirradiated oxide-particle number density 7.0 × 1022 m−3, dispersion distance 91 nm; after 101 dpa, number density 8.0 × 1022 m−3, dispersion distance 90 nm; average oxide size remained <10 nm	At 101 dpa, swelling remained <1%; reported values 0.45% for PM1000e and 0.25% for PM1000e + 0.5Hf	Oxide particles remain densely dispersed after irradiation and trap He/vacancies at particle–matrix interfaces rather than grain boundaries	117
Ferritic ODS steel under irradiation	MA957	Average Y–Ti–O cluster number densities measured between 2 and 6 × 1024 m−3; segregation of Cr-rich features at 412 °C–109 dpa had number density 9 × 1023 m−3 and Guinier radius 2.1 ± 0.1 nm	Irradiation conditions: 412 °C–109 dpa, 550 °C–113 dpa, 670 °C–110 dpa	High density of nanofeatures remains after high-dose neutron irradiation, with temperature-dependent density reduction and evidence of ballistic dissolution/reformation	118

Values are representative literature data selected to illustrate microstructure–property relationships across different ODS alloy systems. Reported values depend on alloy composition, processing conditions, and testing environment; therefore, the table is intended for comparative purposes rather than exhaustive coverage.

Influence of alloying elements

Miller et al. ¹³⁸ studied the influence of Cr and W on the thermal stability of 12 and 14Cr ODS steels. They observed limited coarsening after thermal exposure at 1300 °C for 24 h, irrespective of Cr content, owing to the slow diffusion of Ti and Y₂O₃ from smaller particles. More coarsening in the alloy without W is associated with the higher free energy of W-O interaction. Ratti et al. ¹³⁹ studied the influence of Ti addition on the precipitation and coarsening of Y₂O₃ particles during annealing of Fe-18Cr-1W and Fe–18Cr–1W-0.8Ti (wt.%). The mean radius of Y₂O₃ particles after thermal exposure at different temperatures for both the alloys is shown in Figure 6. They observed superior resistance against coarsening in the alloy containing Ti at all temperatures. Later, Yan et al. ¹⁴⁰ studied the influence of the Hf addition on the thermal stability of 16Cr ODS steels.

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Figure 6.

Influence of Ti addition on nanoparticle coarsening. Reproduced from Ref. ¹³⁹

The experiments were conducted at different temperatures ranging from 750 to 1150 °C for 100 h. The addition of Hf has refined particle size in the as-received condition and leads to less coarsening under extreme thermal conditions. The 16Cr ODS steel without Hf exhibits minimal coarsening at 750 °C, followed by a bimodal particle-size distribution at 950 and 1150 °C. They observed that the specimens heat-treated at 950 °C contain approximately 25% of particles larger than 20 nm. This number further increases at 1050 °C, with nearly 70% of particles exceeding 20 nm in size. However, particle coarsening is limited to 10% and 13% at 950 °C and 1050 °C, respectively. ¹³⁹ Recently, Yang et al. ¹⁴¹ studied the influence of Ce on the thermal stability of FeCrAl ODS steel. They observed that higher Ce content resulted in a higher density of oxide particles, such as Y-Ce-O, Ce-Ti-O, and Ce-Al-O. However, these specimens contained fewer oxide particles, such as Y-Ti-O, Y-Al-O, and M-O, due to the increased Ce addition stabilizing the Ce-based nanoparticles.

Additionally, the size of oxide particles is almost 30% lower in the specimen containing higher Ce content than in those with lower Ce content. The grain size of the specimen with higher Ce content is finer due to the uniform distribution of nanoparticles, resulting in greater resistance to grain-boundary sliding. In this context, the thermal stability of ODS steels is significantly influenced by alloying elements like W, Ti, Hf, and Ce, which enhance resistance to coarsening.

Short-term annealing

Oksiuta et al. ¹⁴² studied the thermal stability of 14Cr ODS steel at 650–1350 °C. They observed no change in yttrium particle size below 1250 °C, and a slight increase from 2.6 to 2.9 nm at 1350 °C. They emphasized that the activation energies of substitutional and interstitial atoms primarily govern nanoparticle coarsening. At lower temperatures, the externally supplied thermal energy is insufficient to drive the diffusion of substitutional atoms such as Y and Ti, thereby reducing coarsening. Zhao et al. ¹⁴³ performed short-term annealing of 14Cr ODS steel at 800, 1000, and 1200 °C for 5 h. The alloy exhibits a mixture of recrystallized and deformed grains in both the as-received condition and after annealing at 800 and 1000 °C. The alloy's tendency to recrystallize at annealing temperatures of 800 and 1000 °C was associated with the presence of fine precipitates that impede recrystallization. The number density of voids increases at 1000 and 1200 °C, accompanied by a fully recrystallized microstructure at 1200 °C. These voids were formed due to the accumulation of micropores obtained during consolidation and HIP. At high temperatures, the pinning effect, in which nanoprecipitates trap micropores, weakens, allowing easier motion of micropores. ¹⁴⁴ The as-hot isostatic pressed specimen contains coarse M₂₃C₆ particles and fine TiC particles. The TiC particles, which have nucleated from defects, are distributed along the grain boundaries where these defects were primarily located. The M₂₃C₆ phase progressively decreases with increasing annealing temperature, thereby facilitating the formation of new TiC particles and causing slight coarsening. They further performed XRD analysis to examine the evolution of secondary-phase fractions across different annealing temperatures, revealing the presence of Y₂Ti₂O₇, M₂₃C_6, and TiC particles. The relative volume fraction of Y₂Ti₂O₇, measured with respect to the total secondary phase particles volume, increased significantly from approximately 0% in the as-HIPed specimen to more than 75% in the specimen annealed at 1200 °C.

Oksiuta et al. ¹¹³ reported that adding 0.3% Y₂O₃ or 0.5% Fe₂Y in a Fe-14Cr-2W-0.2Ti alloy significantly influences their short-term annealing resistance. These alloys were very stable when annealed below 1150 °C. However, the nanoparticles in Fe₂Y ODS alloys underwent partial coarsening after annealing at 1350 °C for 1 h. TEM-based observations of Fe₂Y ODS alloys shown in Figure 7 revealed that the size of a few nanoclusters did not change, whereas the size of some nanocrystals significantly increased. The nanoparticles in the finer nanoclusters were mainly Y- or Ti-oxides. Interestingly, the crystal structure of the nanoparticles in larger nanoclusters of Fe₂Y ODS alloy was FCC, as shown in the inset of Figure 7(c), which differs from the typical body-centered crystal structures observed in nanoparticles of ODS alloys. The Y₂O₃ ODS alloy exhibited greater resistance to coarsening following short-term annealing.

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Figure 7.

(a) and (b) ODS Fe₂Y, nanoclusters, and (c) ODS Fe₂Y, coarse Y₂O₃ nanocluster and SED pattern. Reproduced from Ref. ¹¹³

As discussed earlier, tempering treatment is employed to reduce brittleness and enhance toughness in the martensitic ODS alloy. Zhou et al. ⁴⁷ performed tempering treatment at 800 °C for 1 h on 9Cr-Hf and 9Cr-Zr ODS steels and found similar sizes of M₂₃C₆ precipitates near martensitic lath boundaries and ferrite boundaries. They also observed that the size of oxide particles has increased by approximately 45% in the 9Cr-Zr steel. In contrast, the particle size remains unchanged in the 9Cr-Hf steel, revealing better thermal stability of Y-Hf-O particles. Thus, the short-term annealing resistance of ODS alloys strongly depends on composition, with Y₂O₃-based alloys exhibiting superior coarsening resistance, and Hf additions enhancing nanoparticle stability. High-temperature annealing, even for short durations, promotes recrystallization and phase evolution, which may affect the mechanical properties.

Long-term annealing and precipitate characteristics

Mao et al. ¹⁴⁵ conducted a comprehensive study on the thermal stability of 12Cr ODS steel produced by MA and HIP. The alloy was subjected to thermal exposure at 1250 °C for 500 h. After HIP, orthorhombic YCrO₃ and monoclinic YTaO₄ were formed. They observed that the monoclinic YTaO₄ structure formed during HIP transformed into an FCC Y₃TaO₇ structure during annealing, attributable to reduced oxygen solubility and the dissolution of fine particles. They suggested that the ODS alloy be processed at temperatures below 1200 °C. However, the role of annealing time was not considered in the research, even though it generally plays a crucial role in the coarsening of these particles. Oksiuta et al., ¹⁴⁶ studying 14Cr ODS steel, investigated the effect of annealing time on the thermal stability of both grains and particles. They observed that the grains do not recrystallize during short annealing at temperatures up to 1250 °C. However, the larger grains coarsen due to recrystallization as the temperature exceeds 1250 °C. The change in grain size was limited by the pinning effect of the oxide particles, even at temperatures above 1250 °C. Long-term annealing at 750 °C for 10000 h does not change the grain size. However, the fraction of larger oxide particles increased as new Al-rich oxide particles precipitated, compromising alloy properties. Xu et al. ¹⁴⁷ aged 12, 16, and 18Cr ODS steels at 480 °C for 2000 h. After aging, they observed no grain- or oxide-particle coarsening. They also observed carbide formation, such as M₂₃C₆ (cubic crystal structure) in 12Cr ODS and M₂C (hexagonal crystal structure) in 16Cr ODS steel. They noted that carbide particle coarsening is generally observed in non-ODS alloys, such as P92 steel. However, no significant change in the carbide particle size was evident in the ODS alloys. They attributed this peculiar behavior to the presence of dispersed oxide particles. However, a more detailed analysis could have been performed to determine why the carbide particles do not coarsen in ODS alloys.

Zilnyk et al. ¹⁴⁸ studied the stability of nanoclusters in 9Cr ODS steels exposed at 800 °C for 4320 h. They evaluated the change in particle size using APT and found no change before and after thermal exposure. They concluded that solute atoms become trapped within V- Cr nanoclusters, delaying the precipitation of intermetallic compounds and thereby degrading mechanical properties. Zheng et al. ¹⁴⁹ studied the long-term thermal stability of 9Cr ODS steel at 700 °C for 10000 h. The grain morphology remains similar, showing no recrystallization after the aging treatment. They also emphasized that complete dislocation recovery was not observed in the microstructure, owing to the pinning effects of the oxide particles. The alloy contained only carbides distributed along boundaries and Y-Ti-O oxides. The carbides coarsened with increasing thermal exposure time. However, the size and shape of the Y-Ti-O particles did not vary much after aging, as shown in Figure 8. Gao et al. ¹⁵⁰ studied the thermal stability of 9Cr ODS steel aged for 5000 h at 600 °C.

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Figure 8.

Evolution of Y-Ti-O oxides and carbides at: (a, a’) before aging, (b,b’) 700 °C/100 h, (c,c’) 700 °C/5000 h, and (d,d’) 700 °C/10000 h, respectively. Reproduced from Ref. Zheng et al. ¹⁴⁹

They observed no significant changes in grain or oxide particle size after aging. The alloy's superior thermal stability was attributed to the coherency between the nanoparticles and the matrix, which reduces interfacial energy. Additionally, the lower Cr content in the alloy minimized the formation of the Cr-rich α’ phase, which is associated with embrittlement and early failure. In most cases, ODS steels demonstrate exceptional long-term thermal stability, with minimal grain coarsening and oxide particle growth after 10,000 h of annealing at 600–750°C. In several studies, the grain morphology remains unchanged up to 1250°C, while oxide particles exhibit limited coarsening. Finely dispersed oxide particles effectively inhibit recrystallization, dislocation recovery, and carbide coarsening, thereby maintaining nanoparticle size under prolonged thermal exposure.

Thermal cycling

Oñoro et al. ¹⁵¹ studied the influence of thermal cycling between 400 and 800 °C on the thermal stability of 14Cr ODS steel. They performed microstructural analyses on three different specimens. The first specimen was annealed at 1000 °C for 2 h, while the second and third specimens were obtained by subjecting the first specimen to thermal cycling between 400 °C and 800 °C, followed by aging at 600 °C for 2000 h, respectively. They observed coarse Cr-W-rich and fine Ti-rich secondary phases. EDS analysis showed that the Cr-W-rich precipitate in the aged specimen had a significant Cr deficiency attributed to Cr diffusion into the matrix.

In contrast, W accumulated in the precipitate due to diffusion from the matrix. The median size of the Cr-W-rich secondary phase was lowest for the annealed specimen and highest for the specimen subjected to thermal cycling. However, the highest area fraction of secondary phases was observed in the aged specimen due to easier W diffusion.

Joining-induced microstructural stability

In addition to thermal exposure during service, microstructural stability in ODS alloys can also be significantly affected during joining and fabrication processes. Conventional fusion welding methods are generally unsuitable for ODS alloys because melting and resolidification disrupt the finely dispersed oxide population. During the liquid phase, oxide particles may coarsen, agglomerate, or redistribute, leading to a non-uniform dispersion and degradation of high-temperature mechanical properties in the weld region.

To overcome these limitations, solid-state joining techniques have been explored. Diffusion bonding has been demonstrated as an effective approach for joining ODS alloys while preserving nanoscale oxide dispersion. For example, Noh et al. showed that high-Cr ferritic ODS steel could be successfully joined by diffusion bonding at elevated temperatures without forming voids or inclusions, and with minimal degradation of the oxide particle distribution across the bonding interface. ¹⁵² Similarly, friction stir welding (FSW) has been identified as a promising technique for ODS alloys, as it enables joining through severe plastic deformation in the solid state without melting. Chen et al. reported that friction stir welding of PM2000-type ODS alloys produced a stirred region with fine, equiaxed grains and a relatively uniform distribution of Y–Al–O particles, with an average particle size of approximately 21 nm after recrystallization treatment. ¹⁵³ These observations indicate that maintaining dispersoid stability during joining is critically dependent on avoiding melting and controlling microstructural evolution in the solid state. Although solid-state techniques can preserve oxide dispersion and mechanical integrity, challenges remain in optimizing joint microstructure, controlling recrystallization behavior, and ensuring long-term creep and fatigue performance of the joint region.

Austenitic ODS steels

Sun et al. ¹⁵⁴ studied the influence of high-temperature aging of austenitic ODS steel (20Cr25Ni) at 700 °C for 1000 h. They observed a negligible change in the σ-FeCr phase fraction, in contrast to earlier studies on 25Cr20Ni ODS steel, and suggested that the reduced presence of the σ phase may improve mechanical performance, as this phase is generally considered detrimental. In addition, grain size remained relatively stable as aging time increased. The oxide particles were identified as Y₂Ti₂O₇ with a cubic crystal structure, and their size remained largely unchanged for aging times up to 100 h. However, after 1000 h, a bimodal particle size distribution was observed, with larger particles preferentially located at grain boundaries. This suggests that long-term exposure may promote localized coarsening or redistribution of dispersoids, although the underlying mechanisms require further investigation.

The thermal stability of additively manufactured austenitic 304L ODS steel produced via LPBF was also evaluated after aging at 1000, 1100, and 1200 °C. ¹⁰⁴ The alloy exhibited good thermal stability up to 1100 °C, with retention of the cellular substructure attributed to the presence of oxide nanoparticles. However, at 1200 °C, the size of Y–Si–O nanoparticles increased from approximately 25 to 50 nm after 100 h due to Ostwald ripening, accompanied by changes in local chemistry, including a decrease in oxygen content and an increase in yttrium concentration. These observations indicate that, although nanoscale dispersoids can stabilize fine microstructural features, their coarsening and compositional evolution at sufficiently high temperatures can alter the strengthening response.

Austenitic ODS alloys differ fundamentally from ferritic systems in their thermal evolution because the fcc matrix promotes higher diffusion rates and more pronounced recovery processes, which influence both grain stability and the persistence of strengthening mechanisms. In ferritic ODS steels, the combination of a bcc matrix and a high density of stable nanoclusters is generally effective in suppressing grain-boundary migration and maintaining microstructural stability during long-term exposure. In contrast, in austenitic ODS alloys, grain coarsening is only one aspect of thermal degradation; thermally activated recovery, dislocation rearrangement, and particle coarsening can collectively reduce the effectiveness of dispersoid strengthening at elevated temperatures. As a result, the thermal stability of austenitic ODS alloys must be evaluated not only in terms of grain growth, but also in terms of how matrix structure influences recovery kinetics, particle stability, and the retention of nanoscale strengthening features during prolonged service.

Applications of ODS in nuclear engineering

ODS alloys are strong candidates for nuclear reactor applications owing to their exceptional high-temperature strength, superior creep resistance, and enhanced tolerance to irradiation damage. These properties originate from their distinctive microstructure, in which a high number density of stable nanoscale oxide dispersoids stabilizes grain boundaries, suppresses recrystallization and grain growth, and impedes dislocation motion. In addition to their strengthening role, dispersoid–matrix interfaces act as efficient sinks for irradiation-induced point defects, thereby mitigating defect accumulation and reducing void swelling. Void swelling in structural steels arises from the incomplete recombination of irradiation-generated interstitials and vacancies. At low doses, partial recombination provides a limited self-healing capability. With increasing dose, however, the formation and growth of dislocations and dislocation loops—biased sinks for interstitials—drive a net interstitial flux toward these defects, while vacancies preferentially migrate toward voids. This imbalance promotes void nucleation and growth, leading to volumetric swelling and progressive degradation of mechanical integrity under reactor operating conditions. ¹⁵⁵ By increasing defect sink density and modifying defect transport pathways, ODS alloys significantly suppress swelling relative to conventional steels. ODS alloys considered for nuclear applications can be broadly classified into three categories. ¹⁵⁶

I. Ferritic–martensitic ODS alloys, including Fe–Cr–based systems such as Fe-9Cr, Fe-12Cr, and Fe-14Cr, exhibit excellent resistance to void swelling and irradiation-induced creep. Their BCC crystal structure, combined with high dispersoid densities, provides strong defect recombination capability, making them particularly suitable for high-dose fast-reactor and fusion environments.

Optimized ODS for reactor applications

The current first-generation candidate swelling-resistant alloys envisioned to replace austenitics are ferritic and ferritic-martensitic steels, with second-generation alloys being dispersion-hardened variants of these steels. It is well known that, as a class, ferritic alloys swell and creep less than do austenitic alloys.^158,159

Figure 9(a) and (b) shows the exceptional swelling resistance of the MA957 ODS ferritic steel under high-dose irradiation. ¹⁶⁰ As shown in Figure 9(a), even after irradiation at 450 °C to 400 displacement per atom (dpa), the alloy preserves a stable microstructure with negligible void formation, demonstrating strong resistance to radiation-induced damage. The swelling behavior shown in Figure 9(b) further confirms that MA957 exhibits substantially lower volumetric swelling than conventional ferritic–martensitic alloys such as EP-450 and HT-9. This enhanced stability is attributed to the high density of finely dispersed, thermally stable Y–Ti–O nanoparticles, which act as efficient defect sinks and suppress void nucleation and growth. Collectively, these results highlight the capability of ODS ferritic steels to maintain microstructural integrity and dimensional stability under extreme irradiation environments.

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Figure 9.

(a) Microstructure of MA957 ODS steel tube after irradiation at 450 °C to 400 dpa, showing a stable nanostructure with minimal void formation. (b) Dose dependence of swelling in MA957 compared with EP-450 ferritic and HT-9 ferritic–martensitic alloys, illustrating the superior irradiation-resistance and low swelling rate (∼0.2%/dpa) of MA957 at 450 °C. Reproduced from Ref. ¹⁶⁰

There has been significant progress in ODS alloy development for reactor applications; however, transformable 9Cr alloys remain of particular interest among austenite-free variants, as their microstructure can be precisely engineered through controlled α–γ phase transformation. ¹⁶¹ The introduction of excess oxygen plays a critical role in this process by (1) promoting the formation of fine Y–Ti complex oxides and (2) influencing the matrix carbon activity, thereby modifying the free-energy difference between the α and γ phases. As a result, the α/γ interface motion is effectively pinned by the oxide particles, weakening the driving force for transformation and stabilizing the nonequilibrium residual ferrite phase. Figure 10(a) and (b) presents transmission electron micrographs of oxide dispersions within the residual ferrite and tempered martensite of a dual-phase 9Cr-ODS ferritic steel. ¹⁶¹ Both regions contain a high density of nanoscale Y–Ti–O particles, with those in the residual ferrite being slightly finer and more uniformly distributed. This refined dispersion, coupled with the nanocrystalline substructure retained from mechanical alloying, accounts for the markedly higher yield strength of the residual ferrite (∼1360 MPa) compared with the tempered martensite (∼930 MPa).

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Figure 10.

Transmission electron micrograph of (a) oxide particles in the residual ferrite and (b) in the tempered martensite of a dual-phase 9Cr-ODS ferritic steel. The residual ferrite contains slightly finer, more uniformly distributed Y–Ti–O nanoparticles than the tempered martensite, contributing to its higher strength and enhanced creep resistance. Reproduced from Ref. ¹⁶¹

The resulting dual-phase architecture behaves as a composite microstructure in which the hard residual ferrite enhances creep strength. At the same time, the tempered martensite provides ductility and damage tolerance under high-temperature loading. Previous studies have shown that adding Ti can reduce oxide particle size by approximately one order of magnitude in both Cr-ODS ferritic steel and 9Cr-ODS martensitic steel. ¹⁶² Repeated impacts during MA destroy Y₂O₃, introducing Y and O atoms into the matrix. Next, due to the very low solubility of Ti in the Fe-Cr binary system, Ti interacts with Y and O to first form metastable Y₂TiO₅ (Hexagonal) and then stable Y₂Ti₂O₇ (cubic). The Y₂Ti₂O₇ phase is favored since it has a lower interface energy than Y₂O₃.^54,162–164

The interfacial structure and crystallographic coherency of Y–Al complex oxides play a decisive role in determining the mechanical stability of Al-alloyed ODS steels. Figure 11(a) and (d) shows high-resolution TEM images and corresponding FFT patterns of the nanometer-scale oxides identified in Fe–15.5Cr–2W–0.1Ti–4Al–0.35Y₂O₃ steel extruded at different temperatures. ¹⁶³ Two distinct polymorphs are observed: semi-coherent YAP (orthorhombic YAlO₃, perovskite structure) particles in the specimen extruded at 1150 °C (SOC-9) and coherent YAH (hexagonal YAlO₃) particles in the specimen extruded at 1050 °C (SOC-9-ET-1). The semi-coherent YAP particles exhibit a lattice misfit of approximately 12%, whereas the coherent YAH particles maintain nearly complete lattice matching with the ferritic matrix.

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Figure 11.

(a) HRTEM micrograph of a semi-coherent YAP (orthorhombic YAlO₃) particle with surrounding matrix in SOC-9 and (b) corresponding FFT pattern. (c) HRTEM micrograph of a coherent YAH (hexagonal YAlO₃) particle with surrounding matrix in SOC-9-ET-1 and (d) corresponding FFT pattern. The YAP particle exhibits a semi-coherent interface with a lattice misfit of ∼12%. In contrast, the YAH particle maintains nearly full coherency with the ferritic matrix, reflecting a transition from semi-coherent to coherent oxide structures as the extrusion temperature decreases. Reproduced from Ref. ¹⁶³

Reducing the extrusion temperature refines the oxide size and promotes a phase and coherency transition from semi-coherent YAP to coherent YAH, thereby enhancing thermodynamic stability and reducing interfacial energy. Consequently, samples containing predominantly coherent YAH particles exhibit superior high-temperature strength and creep resistance due to improved particle–matrix bonding and enhanced defect trapping at coherent interfaces.

Issues and challenges in the radiation environment

The stability of dispersoids under irradiation varies for different ODS materials. There is no universal trend. The dispersoid behavior also varies with the type of irradiation (proton, heavy-ion, or neutron) and is strongly dependent on irradiation temperature. The evolution of dispersoids is closely coupled with other microstructural changes, such as void swelling and dislocation development. The following behaviors have been reported.

Yttria dispersions at extremely small sizes may be unstable due to ballistic dissolution and/or amorphization when the interfacial energy becomes significant at high interfacial-to-volume ratios. ¹⁶⁵ Since interfacial energies in ferrite and martensite phases differ, the critical defect density required for amorphization is different in dual-phase ODS alloys. Therefore, the amorphization resistance of yttria dispersions may be different from traditional ODS alloy types.

Although tempered martensite grains exhibit superior swelling resistance, this intrinsic property may change if oxide particles within martensite induce strongly biased defect trapping. A much stronger interstitial trapping at the interface will lead to grain-void swelling due to reduced defect recombination. In comparison, stronger vacancy trapping will lead to the formation of voids at the dispersoid interface and reduce the incubation time for void swelling. Ideally, the interface should serve as a trap for both interstitials and vacancies, facilitating effective defect annihilation. As an example of “bad” ODS, ¹⁶⁶ Figure 12 (a)-(d) compares unirradiated and Fe-ion irradiated MA956 up to 60, 120, and 180 dpa, with ions entering from the left. Voids are found to nucleate exclusively at the particle-matrix interface. The lines superimposed on Figure 12(d) represent the calculated dpa and injected Fe ions as a function of depth. Furthermore, the large particles are surrounded by a cloud of small precipitates, most likely resulting from recoil dissociation at the particle interface and subsequent reprecipitation in the nearby matrix.

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Figure 12.

Conventional bright field images of (a) the unirradiated, (b-d) irradiated specimens with 60, 120, and 180 peak dpa, respectively, showing that essentially all voids are associated with oxide particles. The insert in (b) shows a void formed on the side of an oxide particle in the 60-peak dpa sample. Note that the insert also shows a cloud of small precipitates surrounding the particle that was probably formed by recoil dissolution and reprecipitation. The insert in (c) shows the irradiation-induced dislocation network at a depth of 1000 nm in the 120-peak dpa sample. The lines superimposed on (d) are the calculated dpa and injected interstitial profiles as a function of depth. Reproduced from Ref. ¹⁶⁶

Beyond ferritic and ferritic–martensitic ODS systems, these challenges are more complex in Ni-based alloys, where irradiation-induced degradation is often governed by helium embrittlement, grain boundary instability, and irradiation-assisted phase evolution rather than by swelling alone. In conventional Ni-based alloys, neutron irradiation promotes helium accumulation at grain boundaries through transmutation reactions, leading to intergranular embrittlement and loss of ductility even at relatively low swelling levels. In addition, irradiation-induced solute segregation and precipitation at grain boundaries further degrade mechanical integrity by facilitating crack initiation and propagation under stress. ¹⁶⁷ The introduction of oxide dispersions in Ni-based ODS alloys provides a potential pathway to mitigate these effects by modifying defect transport and helium partitioning. For example, Oono et al. demonstrated that in PM1000-type Ni-based ODS alloys irradiated to ∼101 dpa, oxide particles remain stable with number densities on the order of ∼10^22–10^23 m⁻³, and more than 80% of cavities are associated with oxide particle–matrix interfaces. In contrast, less than 10% are located at grain boundaries. This redistribution of helium-induced cavities away from grain boundaries results in swelling levels below ∼1% and significantly reduces the likelihood of intergranular embrittlement. ¹¹⁷

However, the presence of dispersoids does not eliminate irradiation-induced degradation mechanisms in Ni-based systems. Instead, it shifts the dominant damage processes toward the particle–matrix interface, where cavity nucleation, local decohesion, and irradiation-assisted chemical evolution may occur. Recent studies on NiMoCr–Y₂O₃ alloys have shown that Y₂O₃ particles can effectively trap helium and reduce bubble volume fraction compared to conventional Ni-based alloys such as Hastelloy N, thereby suppressing irradiation-enhanced chromium diffusion and improving corrosion resistance. At the same time, irradiation-induced vacancies at the Y₂O₃–matrix interface can promote interfacial reactions, leading to the formation of complex oxide structures such as Y₂O₃, YCrO₃, and Cr₂O₃, which indicates that the interface itself evolves chemically under service conditions. ¹⁶⁸ Additional irradiation studies on Ni-based alloys further reveal that high-dose irradiation can drive microstructural instability through mechanisms such as irradiation-induced carbide precipitation, interfacial pitting, serrated grain boundary formation, and non-monotonic surface roughening, particularly at elevated temperatures relevant to advanced nuclear systems. ¹⁶⁹

These observations highlight that, in Ni-based ODS alloys, the role of dispersoids must be understood within a broader framework that includes transmutation-driven gas production, interface stability, and irradiation-assisted phase evolution. While oxide dispersions can significantly enhance defect recombination and reduce grain-boundary damage, they may also introduce new pathways for localized damage accumulation if the interface becomes strongly biased or chemically unstable. Therefore, the key challenge is not simply to retain dispersoid populations but to ensure that particle–matrix interfaces maintain balanced defect-trapping behavior and structural integrity under prolonged irradiation. Addressing this challenge will require coordinated control of dispersoid chemistry, interface coherency, and matrix composition, particularly in environments involving coupled irradiation, high temperature, and corrosion, where multiple degradation mechanisms may operate simultaneously.

Dispersoid evolution under irradiation

The long-term irradiation performance of ODS alloys is fundamentally controlled by the stability and evolution of their nanoscale oxide dispersoids. These particles provide the primary mechanisms for strengthening and radiation mitigation by pinning dislocations, stabilizing grain boundaries, and acting as sinks for irradiation-induced point defects. However, under neutron or ion irradiation, dispersoids are not immutable; they may undergo structural and chemical modification, including changes in size, number density, crystallographic structure, and composition.^170,171 Radiation-induced segregation, transmutation effects, and cascade-driven dissolution–reprecipitation processes can further alter particle–matrix interfaces and local chemistry. Because the mechanical integrity and swelling resistance of ODS alloys are directly linked to this nanoscale architecture, understanding dispersoid stability under coupled thermal and irradiation conditions is essential for predicting long-term reliability in service environments.^84,102,172

Despite extensive experimental investigation, reports on irradiation- induced dispersoid evolution in ODS alloys remain highly inconsistent. Some studies describe particle coarsening accompanied by a reduction in number density and compositional modification, whereas others report refinement, partial dissolution, or negligible changes under nominally similar conditions. These discrepancies do not necessarily reflect experimental uncertainty; rather, they highlight the complex, multivariate nature of dispersoid stability. Differences in alloy chemistry, initial dispersoid coherency, processing history, irradiation dose rate, temperature, and damage spectrum strongly influence the balance between ballistic dissolution, thermally driven diffusion, and interfacial energetics. As a result, direct comparison between studies is often misleading unless these parameters are explicitly considered. A critical synthesis of the literature requires a mechanistic interpretation rather than a simple aggregation of reported observations.

A useful contrast to ferritic ODS systems is provided by Ni-based ODS alloys, in which the evolution of dispersoids under irradiation is more strongly coupled to interfacial chemical modification. Oono et al. ¹¹⁷ showed that in PM1000-type Ni-based ODS alloys irradiated to approximately 101 dpa, the oxide particles remained present after irradiation, with average sizes still below 10 nm and a number density of about 8.0 × 10^22 m^-3, indicating that the dispersoid population itself can remain comparatively stable even under high irradiation doses. This result suggests that, as in ferritic ODS alloys, oxide nanoparticles in Ni-based systems can retain their basic sink function during irradiation. However, the evolution of these particles cannot be assessed solely by retention of size and number density. In NiMoCr–Y₂O₃ alloys, the Y₂O₃–matrix interface was found to undergo pronounced chemical modification under coupled irradiation–corrosion conditions, leading to the formation of interfacial oxide structures such as Y₂O₃, YCrO₃, and Cr₂O₃. ¹⁶⁸ These observations indicate that dispersoid evolution in Ni-based ODS alloys may proceed not only through coarsening, dissolution, or refinement, but also through irradiation-assisted interfacial reactions that alter local chemistry and potentially modify sink efficiency. In this respect, Ni-based ODS alloys illustrate an important extension of the general dispersoid-stability problem: even when oxide nanoparticles remain physically present, their functional role under irradiation may change because the surrounding interface evolves chemically. More broadly, irradiation studies on conventional Ni-based alloys also show that high-dose exposure can promote carbide precipitation, interfacial degradation, and grain-boundary restructuring, emphasizing that dispersoid stability in Ni-rich matrices must be interpreted in the context of concurrent phase evolution and enhanced diffusion kinetics at elevated temperature.

The variability in reported dispersoid evolution behavior across different studies can therefore be understood in terms of the combined influence of alloy composition, irradiation conditions, and characterization methodology. For example, apparent stability of nanoclusters is often reported in ferritic ODS alloys irradiated at intermediate temperatures, where limited diffusion suppresses coarsening, whereas at higher temperatures or in FCC-based systems, enhanced atomic mobility can promote coarsening or interfacial chemical evolution. Similarly, ballistic dissolution and fragmentation are more commonly observed under heavy-ion irradiation or at high defect production rates. In contrast, neutron irradiation at lower dose rates tends to produce more gradual evolution. In addition, differences in particle–matrix interface structure may lead to either balanced defect recombination or biased defect trapping, resulting in contrasting observations such as void suppression or interface-associated cavity formation. Finally, differences between characterization techniques, particularly TEM and atom probe tomography, can lead to variations in the reported particle size, number density, and composition of nanoscale dispersoids. These factors collectively explain why dispersoid evolution under irradiation is often reported inconsistently in the literature, despite being governed by the same underlying physical mechanisms.

Advanced nanoscale characterization techniques have been instrumental in resolving these issues by enabling direct observation of dispersoid morphology, chemistry, and spatial distribution. STEM provides high-resolution structural and chemical information on oxide nanoparticles in the 1–5 nm size range. High-resolution TEM allows direct imaging of dispersoid crystal structure and morphology ¹⁷³ while selected-area electron diffraction (SAED) enables phase identification and detection of irradiation-induced structural changes. High-angle annular dark-field (HAADF) STEM provides Z-contrast imaging, enabling the distinction of nanoparticles with differing compositions from the surrounding matrix. Complementary spectroscopic techniques such as energy-dispersive X-ray spectroscopy (EDS) and electron energy loss spectroscopy (EELS) provide compositional and oxidation-state information at nanometer-scale spatial resolution, with EELS being particularly valuable for detecting light elements and probing bonding environments. ¹⁷⁴ Energy-filtered TEM (EFTEM) further allows visualization of elemental distributions and compositional heterogeneity across the microstructure. ¹⁷⁵ APT has emerged as a particularly powerful technique for characterizing ODS alloys, as it enables three-dimensional atomic-scale mapping of nanoclusters with near–parts-per-million chemical sensitivity.^{118,138,148176–178} By reconstructing the spatial and chemical distributions of individual atoms, APT provides quantitative information on dispersoid size, number density, composition, and solute segregation at particle–matrix and grain-boundary interfaces.^179,180 This capability is especially important for assessing irradiation-induced elemental redistribution and validating the chemical stability of oxide nanoclusters, which cannot be fully resolved by projection-based microscopy techniques alone.

Figure 13 shows representative APT results for Y–Ti–O nanoclusters in MA957 ODS ferritic steel following high-dose neutron irradiation (∼100–110 dpa) at different temperatures. ¹¹⁸ At lower irradiation temperatures (≈412 °C), the reconstructed ion maps reveal a dense population of fine nanoclusters. In contrast, increasing the irradiation temperature leads to a progressive decrease in the number density and to the coarsening of the remaining clusters. Quantitative analysis shows that the cluster number density decreases from approximately 5 × 10²⁴ m⁻³ at 412 °C to below 2 × 10²⁴ m⁻³ at 670 °C. Notably, the average Y–Ti–O stoichiometry remains nearly unchanged across this temperature range, with Ti:O ≈ 1 and Y: Ti < 1. These observations indicate that irradiation primarily modifies the population and spatial distribution of nanoclusters rather than their intrinsic chemistry, supporting a ballistic dissolution–reformation mechanism with strong temperature dependence.

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Figure 13.

Atom probe tomography analysis of high dose MA957 at selected irradiation temperatures (a) Ion maps of Y–Ti–O solute ions in 412 °C–109 dpa specimen. (b) Plotted above are the trends in cluster density as a function of irradiation temperature. For each condition, a box and whisker plot indicates the character of the distribution, for which an average and standard deviation were calculated. The error bars associated with the averages are too small to be visible on this scale. (c) Plotted above are the composition distributions of the Y–Ti–O clusters as a function of irradiation temperature. For each element and temperature condition, a box and whisker plot, as well as an average and standard deviation, characterize the compositional distribution. Reproduced from Ref. ¹¹⁸

The enhanced diffusion at higher temperatures promotes coarsening and reduces cluster density. Taken together, these findings demonstrate that a single universal trend cannot describe the evolution of dispersoids in ODS alloys. Instead, it reflects a competition between irradiation-induced dissolution, thermally activated diffusion, and interfacial stability, all of which are strongly influenced by alloy composition and processing history. Consequently, robust interpretation of microstructural evolution requires correlative STEM–APT analysis combined with mechanistic frameworks that explicitly account for these competing effects. Such an approach provides the foundation for understanding how microstructural stability governs the mechanical and irradiation performance of ODS alloys, as discussed in the following sections.

CALPHAD-based thermodynamic modeling has been used to assess phase stability, oxygen solubility limits, and the driving forces for complex oxide formation (e.g., Y–Ti–O and Y–Zr–O), thereby guiding design of alloy chemistry and processing windows that suppress undesirable coarse oxides. Phase-field simulations have been employed to model nanoparticle nucleation, growth, and coarsening kinetics, offering mechanistic insight into dispersoid stability during thermal exposure and irradiation. More recently, data-driven and machine-learning approaches have emerged as tools to accelerate ODS alloy development by correlating processing parameters (milling energy, consolidation route, thermal history) with microstructural descriptors and mechanical performance.^181,182 These approaches are particularly promising for navigating the high-dimensional compositional and processing space associated with ODS systems, where purely empirical optimization is inefficient. While still at an early stage, the integration of CALPHAD, phase-field modeling, and machine learning with high-resolution experimental characterization represents a powerful pathway toward predictive, mechanism-guided ODS alloy design. Dispersoids under irradiation can shrink or dissolve, coarsen or grow, undergo amorphization, or remain essentially unchanged. ¹⁷¹ The stability of dispersoids is believed to be determined by their size, density, interface coherency, and irradiation conditions. It also depends on the matrix phases. Although various behaviors have been reported, most ODS alloys exhibit size shrinkage after irradiation. To interpret these diverse experimental observations, rate-theory and cluster-dynamics–based modeling approaches have been widely employed to describe irradiation-driven dispersoid evolution. These models treat changes in dispersoid size as the result of competing thermally activated recovery processes and irradiation-induced dissolution driven by displacement cascades, providing a quantitative framework for rationalizing experimentally observed shrinkage, coarsening, or stabilization trends. Figure 14 shows the size evolution in ion-irradiated 12Cr ODS. ¹⁸³ Due to the limited penetration depth of ions, irradiation damage is confined to the near-surface region, approximately 1.4 microns for 3.5 MeV Fe ions. Within the damaged region, dispersoids shrink.

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Figure 14.

Mean dispersoid diameter in the tempered martensite phase as a function of depth, irradiated at different temperatures, with error bars being the standard deviations of experimental measurements. Due to potential interference from irradiation-induced defect clusters, the reported data at 325 °C represent the upper limit of dispersoid size. Reproduced from Ref. ODS. ¹⁸³

Additionally, the final dispersoid sizes are clearly temperature-dependent: the higher the irradiation temperature, the larger the size. The size and coherency of a dispersoid were found to be correlated in different ODS alloys. Therefore, any size prediction must consider the role of coherency, as it determines the free energy of a dispersoid and the pathways along which free energy is minimized. The free energy of a dispersoid is a combination of interfacial energy and elastic energy.

Small dispersoids prefer to remain coherent to minimize surface energy. As dispersoids grow, elastic energy becomes dominant over interfacial energy, causing particles to coarsen and adopt incoherent interfaces. A semi-coherent state is an intermediate condition in which dislocations form at the interface to relieve strain energy partially. From a modeling perspective, this competition is commonly captured using diffusion-based rate equations that couple solute transport, interfacial energetics, and the production of irradiation damage. Such formulations are not intended to predict absolute dispersoid sizes but rather to identify stability regimes and dominant mechanisms under given temperature and dose-rate conditions.

The competition between recovery and irradiation-induced dissolution determines the changes in size induced by irradiation. Recovery or growth is driven by thermal diffusion and can be approximated by the following Equation (1) ¹⁸⁴ :

d r d t = D r ⋅ c − c r c p − c r , c r = c ∞ exp ( 2 γ i υ a t k T r )

Equation (1): Recovery or growth driven by thermal diffusionWhere d r d t is the radius change rate, D, D is the solute diffusion coefficient, c is the solute concentration in the matrix, c r is the solute concentration at the interface, c p is the solute concentration in the dispersoid, c ∞ is the solubility limit at a flat interface, γ i is the unit interfacial energy at the dispersoid-matrix interface, υ a t is the average atomic volume in the dispersoid, T is the temperature, and k is the Boltzmann constant. Under irradiation, damage cascades dissolve dispersoids. The competition between these processes determines the final equilibrium size, as the solution of the following Equation (2) ¹⁸⁵ :

d r d t = D r ⋅ c − c r c p − c r − K ψ = 0

Equation (2): Final equilibrium size.where K is the dpa rate, and ψ is a parameter describing the efficiency of damage cascades in dissolving dispersoids. Figure 15 shows the numerous solutions of the above equation, with a combined unit introduced to cover a variety of dispersoids. ¹⁸³ The solid black line represents the stability boundary. Dispersoids initially located within the curve tend to grow and reach a relatively converged size (the top portion of the black line). At the same time, dispersoids outside the curve tend to shrink to the converged size if their original sizes are too large or completely dissolve if their original sizes are small and fall below the black curve. The stability boundary expands to dashed lines as c / c ∞ decreases. Although simplified, these rate-theory–based models have proven effective in capturing first-order trends in dispersoid stability across a wide range of ODS alloys and irradiation conditions. When combined with experimental APT and TEM measurements, they provide mechanistic insight into why certain dispersoid populations converge on stable sizes, whereas others dissolve under irradiation.

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Figure 15.

Stability boundary of dispersoid as a function of combined units. Reproduced from Ref. ¹⁸³

Dispersoid reprecipitation

Under heavy-ion irradiation, large damage cascades may shrink or dissolve dispersoids. If the solute concentration is sufficient, the displaced atoms may re-nucleate into new oxide particles. However, the newly nucleated dispersoids may not retain the same structure or phase as the original dispersoids. Figure 16 schematically illustrates the dissolution and re-nucleation process. ¹⁸⁶

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Figure 16.

Schematics of the oxide dissolution-reprecipitation process in coarse grains of COS-2. Reproduced from Ref. ¹⁸⁶

The Y₂Ti₂O₇ dispersoids evolve into irregular shapes due to ballistic collisions. The reprecipitated dispersions form as Y-Ti-O nanocrystals, which are smaller in size and have a higher density compared to the original dispersoids.

Mechanical property changes after irradiation

Irradiation hardening is a widely observed phenomenon, but its underlying causes vary. The introduction of dislocations, dislocation loops, void swelling, and new precipitates all contribute to hardening. Mechanical property studies on irradiated ODS alloys have largely utilized microscale mechanical testing,^187,188 including indentation, micropillar compression testing, and cantilever bending experiments. For nanoindentation tests, the size effect arising from localized damage near the surface of an ion-irradiated ODS must be appropriately evaluated and accounted for in data analysis. Early studies have shown that the shape of the plastic zone differs between ion-irradiated and neutron-irradiated samples. ¹⁸⁹ Neutron-irradiated specimens exhibit a more spherical plastic zone, indicating extensive plastic flow in the radial direction.

In contrast, ion-irradiated specimens develop an elongated plastic zone that extends farther in the indentation-depth direction than radially. Micropillar compression testing can be applied either on the surface of an irradiated specimen or on the cross-section of the irradiated specimen. The latter has the advantage of increasing the homogeneity of obstacles along the indentation direction. The typical micropillar diameter is a few microns. The size effect has been a subject of intensive research.^190–192 It is well known that in small specimens, dislocation starvation occurs, and the measured yield stress exceeds the bulk value. However, this size effect may be less pronounced in ODS alloys. Figure 17 shows the systematic effect of specimen size in specimens with varying obstacle densities. ¹⁸⁹ The yield stress exceeds the bulk value in the region dominated by the intrinsic size effect (due to dislocation starvation). As pillar size increases, the value enters a transition regime in which specimens with fewer obstacles exhibit a decrease in strength. This occurs because, in these specimens, the size effect is dominated by specimen size rather than by obstacle distance, particularly in samples with fewer obstacles, such as unirradiated specimens or those without dispersoids. Later, the yield strength approaches the bulk value asymptotically. In contrast, for samples with a high obstacle density, such as irradiated samples or alloys containing a high density of dispersoids, the yield strength rapidly approaches the bulk value. These studies are useful for guiding the design and interpretation of compression testing.

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Figure 17.

Illustration of the specimen size effect and the influence of irradiation, with transition dimension indicated by a vertical dashed line. ¹⁸⁹

Challenges and insights in simulating neutron damage

Using ion irradiation to emulate neutron irradiation has been widely utilized to reduce costs and expedite material screening.^193–195 The shortcomings of ion irradiation are well known and documented. Proton irradiation can introduce a relatively deeper region of damage than heavy-ion irradiation, thereby facilitating mechanical testing. Heavy-ion irradiation is the most efficient method for inducing high levels of damage, but caution is warranted. The surface-defect sink property reduces void swelling near the surface. ¹⁹⁶ The additional atoms introduced via ion implantation cause a local imbalance in defect concentrations, resulting in a higher concentration of interstitials near the projected range and reduced void swelling. ¹⁹⁷ These two complexities serve as boundaries that constrain the region from which reliable data can be extracted. Recently, a quantitative method for determining the safe analysis zone has been developed, based on the consistency of local dpa dependence, ¹⁹⁸ thereby increasing the credibility of the data. Comparative studies of different irradiation types have shown that, in Fe-9% Cr ODS, dislocation loop size and density are comparable under proton and neutron irradiation. ¹⁹⁹ However, the oxide size and density are lower in neutron irradiation.

Additionally, the composition of oxides differs in neutron irradiation. In general, protons can reproduce the voids and loop microstructures observed in neutron irradiation at the same damage level and temperature. However, the absence of heavy-damage cascade production and large ballistic collisions results in different dispersoid dissolution under proton irradiation. ¹⁹⁹ These studies highlight the need to revisit nucleation theory and develop a model to predict nucleation sensitivity to irradiation parameters.