Preamble

Research Progress on Environmental Compatibility of Stainless Steels Used in Lead-Cooled Fast Reactors

Jibo Tan¹, Xinrui Zhang¹,², Baoquan Xue¹, Ziyu Zhang¹, Xinqiang Wu¹*

¹CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
²School of Materials Science and Engineering, Northeastern University, Shenyang 110819, China

Abstract

The environmental compatibility of structural materials in liquid lead-bismuth eutectic (LBE) represents a critical bottleneck restricting the development and deployment of lead-cooled fast reactors. While 9–12Cr ferritic/martensitic steels and austenitic stainless steels are preferred candidate materials for key reactor components, they face severe liquid metal corrosion (oxidation and dissolution) in high-temperature LBE environments. Alloying with appropriate amounts of Si or Al is commonly employed to enhance corrosion resistance. This paper reviews recent research progress on the oxidation, dissolution, slow strain rate tensile behavior, creep, fatigue, and crack propagation of 9–12Cr ferritic/martensitic steels, austenitic stainless steels, Si-enhanced variants, and alumina-forming austenitic stainless steels in LBE. The liquid metal embrittlement (LME) sensitivity and damage mechanisms of ferritic/martensitic steels (body-centered cubic) and austenitic stainless steels (face-centered cubic) are discussed, current research limitations are identified, and future research directions are proposed.

Keywords: Lead-cooled fast reactor, Stainless steel, Liquid metal corrosion, Liquid metal embrittlement, Mechanical property

China is actively developing nuclear power in a safe and orderly manner, pursuing a "three-step" strategy of thermal reactors → fast reactors → fusion reactors. While commercial pressurized water reactors currently dominate, advanced reactor types including lead-cooled fast reactors (LFRs), sodium-cooled fast reactors, and thorium-based molten salt reactors are under active development. LFRs employ a closed fuel cycle and can utilize U-238, which has a relative abundance of 99%, compared to less than 1% for U-235 used in PWRs, thereby supporting sustainable development of nuclear fission energy. Pure lead or lead-bismuth eutectic alloy offers advantages including low melting point, high boiling point, favorable neutron economy, and chemical inertness, making it the preferred coolant for LFRs [1–3]. Several LFR designs have been proposed internationally, including ELFR (Europe), BREST-OD-300 and SVBR-100 (Russia), SSTAR (USA), URANUS-40 (Korea), CiADS and CLEAR-I (China), and PBWFR (Japan), with key parameters summarized in [TABLE:1] [3].

LFRs are typically designed for 15–30 year lifetimes. Key components such as fuel cladding and reactor vessels experience coupled damage from irradiation (up to 150 dpa), high temperatures (300–650°C), complex stresses, and dynamic LBE corrosion (at flow velocities of ~2 m/s) during long-term service. Their performance is critical to reactor safety. Consequently, structural materials must meet stringent requirements: (1) excellent comprehensive mechanical properties including high-temperature strength, ductility, creep resistance, fatigue performance, and creep-fatigue interaction; (2) high resistance to irradiation swelling to maintain dimensional stability; (3) good resistance to irradiation hardening and embrittlement; and (4) superior environmental resistance in LBE, including oxidation, dissolution, and LME [2,4].

9–12Cr ferritic/martensitic steels exhibit excellent high-temperature mechanical properties and irradiation resistance, with extensive service data from fossil power plants and irradiation damage data from fusion reactor first-wall applications [5], making them prime candidates for LFR components. Austenitic stainless steels offer good strength, ductility, and corrosion resistance, with substantial operational experience from PWRs, also making them preferred materials. Novel materials such as FeCrAl alloys, high-entropy alloys, and SiC demonstrate outstanding oxidation resistance and are under consideration [2]. Adding Si or Al to ferritic/martensitic and austenitic stainless steels is a primary approach to enhance LBE corrosion resistance. In recent years, Si-enhanced stainless steels have become key materials for fuel cladding and heat exchangers, as evidenced by Russia's BREST-OD-300 construction and reported industrial-scale production of Si-enhanced LBE-resistant stainless steels by the China Institute of Atomic Energy. Therefore, a systematic review of the degradation behavior of 9–12Cr ferritic/martensitic steels, austenitic stainless steels, and their Si-enhanced and Al-containing variants in LBE is warranted.

1. Corrosion Thermodynamic Analysis

Liquid metal corrosion in LBE primarily involves dissolution and oxidation. [FIGURE:1] shows the solubility curves of various metallic elements in LBE, demonstrating that many elements exhibit high solubility at elevated temperatures, particularly Ni, Mn, and Cu [1]. Consequently, Ni-, Mn-, and Cu-rich alloys such as nickel-based superalloys are generally unsuitable for LFRs. However, Ni, Mn, and Cu are typical austenite-stabilizing elements essential for single-phase FCC alloys like 316 stainless steel. To mitigate dissolution corrosion, protective oxide scales must form on the surface to prevent direct contact between the substrate and LBE.

The oxidation of alloying elements in LBE, consuming 1 mol O₂ to form oxide MₓOᵧ under standard atmospheric pressure, follows reaction (1):

$$2x/y M + O_2 = 2/y M_xO_y \tag{1}$$

where M represents an alloying element. Thermodynamic data [6] enable calculation of the Gibbs free energy change (ΔGθ, kJ/mol) for MₓOᵧ formation.

Oxide stability depends thermodynamically on temperature and oxygen partial pressure. The relationships between ΔGθ, enthalpy change (ΔHθ, kJ/mol), entropy change (ΔSθ, J/mol·K), and temperature (T, K) for common oxides in LBE are listed in [TABLE:2]. The dissolved oxygen concentration in LBE follows Sievert's law, relating oxygen partial pressure ($p_{O_2}$, bar) to dissolved oxygen concentration ($C_O$, wt.%) as shown in equation (2) [7]:

$$p_{O_2} = 10^{13.558 - \frac{20162}{T}} \cdot C_O^2 \tag{2}$$

The relationship between ΔGθ and $p_{O_2}$ is given by equations (3) and (4) [7]:

$$\Delta G = \Delta G^\theta + RT \ln(p_{O_2}) \tag{3}$$

$$\Delta G = \Delta H^\theta - T\Delta S^\theta + RT \ln(p_{O_2}) \tag{4}$$

where R is the gas constant (8.3145 J/mol·K). These relationships enable construction of Gibbs free energy-temperature-dissolved oxygen concentration diagrams for oxides in LBE environments, as shown in [FIGURE:2], which discriminates the thermodynamic stability of oxides under various temperature and oxygen concentration conditions. The oxygen affinity of elements in LBE follows the sequence Al > Ti > Si > Cr > Fe > Ni > Pb, indicating that Al, Ti, Si, and Cr oxides remain stable at high temperatures. Therefore, Al, Ti, and Si are commonly added as corrosion-resistant elements to 9–12Cr ferritic/martensitic steels and austenitic stainless steels to enhance LBE corrosion resistance.

Notably, oxidation in LBE also depends on the activity of elements in the alloy, as expressed in equation (5) [8–10]:

$$p_{O_2} = \frac{1}{a_M^{2x/y}} \exp\left(\frac{\Delta G^\theta}{RT}\right) \tag{5}$$

where $a_M$ is the activity of element M (typically < 1). Consequently, the minimum dissolved oxygen concentration required for oxide formation in alloys is generally higher than the values predicted in [FIGURE:2].

2. 9–12Cr Ferritic/Martensitic Steels

2.1 Corrosion Behavior in LBE

9–12Cr ferritic/martensitic steels exhibit excellent neutron irradiation swelling resistance and good high-temperature (<600°C) creep performance, making them prime candidates for LFR fuel cladding. Major grades include T91, P92, HT-9, E911, and Manet II [2,11–13], typically containing ≤0.5 wt.% Si. Extensive studies on their environmental compatibility in LBE reveal that at temperatures below 450°C, dense Fe₃O₄ and FeCr₂O₄ oxide scales form, providing effective protection against oxidation and dissolution. Above 550°C, high-oxygen conditions produce porous outer Fe₃O₄, dense inner FeCr₂O₄, and a Cr-selective internal oxidation zone, while low-oxygen conditions cause dissolution corrosion [13–20]. [FIGURE:3] shows typical oxide cross-sections on T91 steel after 1000 h exposure at 550°C under various oxygen concentrations: oxidation occurs above 1.26×10⁻⁶ wt.% dissolved oxygen, while LBE penetration and dissolution occur below 1.41×10⁻⁸ wt.% [13]. Additionally, high LBE flow velocities (~2 m/s) can erode oxide scales, bringing the substrate into direct contact with LBE and accelerating mass transfer, further degrading oxidation and dissolution resistance [15,16,19].

To improve corrosion resistance, 1–2 wt.% Si is typically added to promote formation of dense Si-rich protective scales. Chemical compositions of typical Si-enhanced ferritic/martensitic steels are listed in [TABLE:3]. Wang et al. [21] reported that Si content must exceed 0.5 wt.% in 9Cr steels to effectively enhance LBE corrosion resistance. M.P. Short et al. [22] studied Fe-12Cr-2Si corrosion at 600–715°C (506 h), finding Fe-Cr-Si-rich scales, particularly nanoscale Si-rich inner oxides, that significantly improved corrosion resistance in both high- and low-oxygen LBE. Rong et al. [23,24] reported that after 10,000 h exposure at 550°C in dynamic oxygen-controlled LBE (0.3 m/s, 10⁻⁷–10⁻⁶ wt.%), 12Cr-Si steel formed Si-Cr-rich oxides at the scale-substrate interface, reducing oxidation rates by approximately 50% compared to HT-9. Pre-oxidation in 1% O₂ atmosphere at 720°C for 1 h further improved corrosion resistance. Polekhina et al. [25] investigated EP-823-sh steel after 2500 h at 630°C in dynamic oxygen-controlled LBE (4–8×10⁻⁷ wt.%, ~2 m/s), observing non-uniform 0.25–18 μm scales composed of Si-Mn-containing Fe-Cr spinel with Si-rich internal oxidation zones. This steel is used for BREST-OD-300 fuel cladding [26,27]. Russia also developed high-Si (1.9 wt.%) EI-852 steel, though it exhibits significant irradiation embrittlement [27]. Yang et al. [28] developed SIMP steel with superior irradiation resistance, high-temperature oxidation resistance, LBE corrosion resistance, and creep performance (at 550°C) compared to T91. Lu et al. [29] developed 9Cr-Si steel, finding Si addition effectively improved corrosion resistance in static oxygen-saturated LBE at 550°C. However, Schroer et al. [30] studied 1.4718 steel at 450°C and 550°C in dynamic oxygen-controlled LBE (10⁻⁷–10⁻⁶ wt.%, 2 m/s, 1200–15,000 h), noting significant improvement over T91 at 450°C but minimal effect at 550°C, where a thin outer spinel and deep Si-selective internal oxidation along grain boundaries were observed. Shi et al. [31] similarly found that Si additions (0.7, 1 wt.%) to 9Cr steel after 2000 h at 550°C and 600°C (10⁻⁶ wt.% O₂) produced thinner outer spinel scales and deeper Si-oxidized internal zones, suggesting altered scale structure improved scale-substrate adhesion and corrosion resistance, as shown in [FIGURE:4].

Irradiation-LBE synergy significantly affects fuel cladding corrosion. Yao et al. [32] studied in-situ Ar⁺ irradiation (1.36 dpa) coupled with dynamic LBE corrosion (350°C, saturated oxygen, 0.6 m/s, 92 h) on SIMP steel (11Cr, 1.43Si), finding irradiation-enhanced diffusion increased scale thickness from 110 nm (0 dpa) to 500 nm (1.36 dpa). Frazer et al. [33] observed similar irradiation-accelerated oxidation on HT-9 steel (12Cr, 0.4Si) under proton irradiation (~22 dpa) in LBE (420°C, saturated oxygen, 80 h), with scale thickness increasing from 1 μm to 13 μm. Demir et al. [34] recently studied Fe-12Cr-2Si under proton irradiation-LBE corrosion (4 h), finding that at 675°C in low-oxygen LBE, Cr-Si-rich oxides formed protective scales with minimal irradiation effect. However, current irradiation-corrosion coupling experiments are limited to <100 h durations and maximum doses of 22 dpa, far from the 15–30 year design life and 150 dpa fuel cladding doses of LFRs, necessitating long-term coupled data for component design and safety assessment.

2.2 Mechanical Behavior in LBE

The mechanical performance of ferritic/martensitic steels in LBE critically affects component safety, encompassing slow strain rate tensile (SSRT), fatigue, creep, and crack propagation. Numerous studies have evaluated LME sensitivity as functions of temperature, dissolved oxygen, strain rate, irradiation hardening, and Si content [35–44]. Results show that elongation-temperature curves typically exhibit a "ductility trough," with minimum ductility and maximum embrittlement around 350°C, characterized by quasi-cleavative fracture features [40,41] as shown in [FIGURE:5]. At low oxygen concentrations, protective oxide formation is inhibited, allowing direct LBE-substrate contact, dissolution corrosion, and wetting-induced quasi-cleavage with significantly reduced ductility. Strain rate affects LME sensitivity by influencing oxide rupture/repair kinetics and LBE-substrate contact time; higher strain rates increase embrittlement in oxygen-saturated LBE, while lower rates are more detrimental in low-oxygen conditions. Irradiation effects depend on hardening: when irradiation and SSRT temperatures are similar (pre-irradiated samples tested in LBE), irradiation hardening enhances LME sensitivity; when irradiation temperature exceeds test temperature, no hardening occurs and LME sensitivity remains unchanged [42]. Although Si improves corrosion resistance, it promotes Laves phase precipitation, raising the ductile-to-brittle transition temperature [24]. First-principles calculations indicate Si significantly reduces surface energy and Young's modulus, increasing BCC Fe embrittlement susceptibility in LBE [44]. However, systematic SSRT studies on Si-enhanced ferritic/martensitic steels are lacking.

Low-cycle fatigue (LCF) performance is influenced by strain amplitude, temperature, dissolved oxygen, and strain rate [45–51]. T91 steel LCF life in LBE increases with decreasing strain amplitude, with environmental effects more pronounced at high strain amplitudes [FIGURE:6a] [45]. Fatigue life-temperature curves show "valley and peak" characteristics: T91 LCF life increases with temperature in 150–250°C and 350–500°C ranges, but decreases in 250–350°C and 500–550°C ranges [46,47]. LCF life in oxygen-saturated LBE exceeds that in low-oxygen LBE at 350°C [FIGURE:6a], though high-temperature (550°C) oxygen effects require further investigation. Strain rate effects are temperature-dependent: minimal impact at 350°C [FIGURE:6c] [45], but significant life reduction with decreasing rate (0.004–0.4% s⁻¹) at 550°C in saturated oxygen [FIGURE:6d] [48]. Current LCF studies focus on low temperatures (≤350°C), necessitating systematic investigation at typical service temperatures (550°C).

Creep, fracture toughness, and crack propagation properties are also degraded. T91, P92, and HT-9 exhibit 50× higher creep rates in 550–650°C LBE compared to air [11,12,37,52]. FeCrAl coatings on T91 improve creep resistance in LBE at 550°C [53]. Fracture toughness decreases by ~30% in 200–355°C LBE with decreasing loading rate [54–58]. Fatigue crack growth rates increase significantly in 150–450°C LBE (saturated oxygen, R=0.1, 0.1 Hz) due to LBE wetting at crack tips, diffusion, adsorption, and quasi-cleavage cracking [59,60].

3. Austenitic Stainless Steels

3.1 Corrosion Behavior in LBE

Austenitic stainless steels are preferred for LFR components due to excellent fracture toughness, corrosion resistance, and intermediate-to-high temperature mechanical properties, with major grades including 316, 304, and 15-15Ti [2,24,61–63]. Studies show that in high-oxygen LBE, austenitic steels form porous outer Fe₃O₄ and dense inner FeCr₂O₄ scales [16,24,61–66]. Below 450°C, these scales provide adequate protection; above 450°C, scales become ineffective, causing severe intergranular oxidation under high oxygen and dissolution corrosion (particularly Ni-selective dissolution leading to ferritization) under low oxygen. [FIGURE:7] shows intergranular oxidation in 316L after 2000 h at 550°C in saturated-oxygen LBE, especially at high-angle grain boundaries [62]. High LBE flow velocities (~2 m/s) can erode scales, promoting Ni dissolution and accelerating corrosion [66].

Si addition improves corrosion resistance by forming dense Si-rich protective scales. Compositions of typical Si-enhanced austenitic stainless steels are listed in [TABLE:4]. Schroer et al. [67] studied 1.4571 steel (1 wt.% Si) after 40,000 h at 550°C in dynamic oxygen-controlled LBE (2 m/s, 10⁻⁹–10⁻⁵ and 10⁻⁶ wt.%). Samples exposed to low oxygen (10⁻⁹–10⁻⁵ wt.%) suffered severe dissolution, while those at 10⁻⁶ wt.% formed protective Cr-Si-rich scales in localized regions. Kurata et al. [68] found that 18Cr-20Ni-5Si formed continuous submicron Si-rich scales after 3000 h at 450°C and 550°C in static LBE, demonstrating excellent corrosion resistance. Roy et al. [69] reported that 18Cr-15Ni-3.7Si and 21Cr-11Ni-1.6Si formed nanoscale oxide films with good corrosion resistance after 1850 h at 520°C in static LBE (10⁻⁹–10⁻⁴ wt.% O₂). Chen et al. [24,70–72] developed Fe-15Cr-9Ni-2Si steel, which after 3000 h at 550°C in static saturated-oxygen LBE formed outer porous Fe₃O₄ and dense inner Fe-Cr spinel containing uniformly distributed nanoscale SiO₂ particles at the scale-substrate interface and within nanovoids [FIGURE:8]. This Si-enhanced steel showed no significant intergranular oxidation, likely due to improved scale densification and hindered diffusion from the nanoscale SiO₂. Wu et al. [24] also reported significantly reduced oxidation rates for Si-enhanced austenitic steels compared to 316H, with no dissolution after 1500 h in dynamic oxygen-controlled LBE at 550°C (0.3 m/s, 10⁻⁶–10⁻⁷ wt.%). However, Kurata et al. [73,74] noted that Si-enhanced austenitic steels (2.4–5.8 wt.% Si) can still experience dissolution in low-oxygen (10⁻⁸ wt.%) LBE during long-term exposure. In summary, moderate Si addition improves oxidation and dissolution resistance in austenitic stainless steels and represents a primary development direction for LFR vessel materials.

Al addition also enhances corrosion resistance. Alumina-forming austenitic (AFA) steels typically contain Fe-(12–35)Ni-(12–20)Cr-(2–6)Al-(0.1–3)Nb. Zhang et al. [75,76] studied Fe-23Ni-15Cr-3Al in 600°C LBE, observing dissolution at 10⁻⁸ wt.% O₂, but formation of ~100 nm Cr-Al-rich scales with continuous Al₂O₃ at the scale-substrate interface at 10⁻⁶ wt.% O₂, providing effective protection. In saturated oxygen, ~5 μm scales formed with outer Fe₃O₄ + minor Cr-Al oxides, intermediate Fe-Cr spinel, and inner Cr-rich + minor Al-rich oxides, though localized dissolution occurred. Cold-work-induced dislocations and grain boundaries enhanced Cr and Al diffusion, promoting rapid formation of protective scales. Shen et al. [77] reported that Fe-14Ni-14Cr-2.5Al formed continuous nanoscale Al-rich scales with excellent corrosion resistance after exposure to dynamic oxygen-controlled LBE at 550°C (1.8 m/s, 5×10⁻⁷–5×10⁻⁶ wt.%). Tsisar et al. [78] investigated Fe-14Cr-2Mn-20Ni-0.5Cu-3Al and Fe-14Cr-5Mn-12Ni-3Cu-2.5Al after 10,000 h at 500°C in oxygen-controlled LBE (10⁻⁹–10⁻⁶ wt.%), finding submicron Al-Cr-rich scales with good corrosion resistance, though the high-Ni steel experienced localized dissolution at low oxygen (10⁻⁹ wt.%). These results demonstrate that AFA steels form protective Al-rich scales, particularly continuous Al₂O₃ at the scale-substrate interface, though localized dissolution risks remain under low-oxygen (<10⁻⁸ wt.%) conditions.

3.2 Mechanical Behavior in LBE

Research on austenitic stainless steel mechanical behavior in LBE includes SSRT, creep, fatigue, and crack propagation. Sapundjiev et al. [79] conducted SSRT on 316L at 200°C (5×10⁻⁶ s⁻¹), finding no embrittlement in unirradiated or irradiated (1.5 dpa) conditions. Bosch et al. [80] and Stergar et al. [81] confirmed no embrittlement in irradiated (1.5 dpa) 316L at 200–450°C. Petersson et al. [82] tested AFA steels in low-oxygen (10⁻¹⁰ wt.%) LBE (140–600°C, 5×10⁻⁵ s⁻¹), observing no embrittlement below 550°C but intergranular cracking above 570°C. Serre et al. [83] also reported reduced elongation and LME in AFA steels in liquid lead above 500°C. Gong et al. [84] studied 15-15Ti creep behavior at 550°C and 600°C in LBE (10⁻¹⁰ wt.%–saturated oxygen), finding significantly increased creep rates due to dissolution corrosion and intergranular cracking.

Yas'kiv et al. [85] investigated Fe–18Cr–10Ni LCF behavior at 350°C in liquid lead, showing decreased fatigue life with increasing strain amplitude. Ding et al. [86] studied 316LN tubular specimens (internally filled with LBE) at 400°C, finding fatigue life reduction only at high strain amplitudes (≥0.8%) compared to air, with no effect at low amplitudes [FIGURE:10]. Kalkhof et al. [87] reported similar LCF life for 316L in 260°C LBE and air at 0.5% strain amplitude. Marmy et al. [88] found comparable 316L fatigue life in oxygen-controlled LBE and air at 300–400°C. Most studies [24,85–88] show fatigue fractures with striation features, indicating ductile fracture and low LME sensitivity. However, Chocholousek et al. [89] observed accelerated crack growth in 316L at 300°C in low-oxygen LBE. Xue et al. [90] studied 316LN fatigue crack propagation at 200–400°C, finding increased rates with temperature and quasi-cleavage features, indicating LME during crack propagation. Current studies focus on low temperatures (≤400°C), requiring systematic investigation at typical service temperatures (400–500°C).

4. Liquid Metal Embrittlement Mechanism

LME significantly degrades mechanical properties (ductility, fracture toughness, creep, fatigue) and threatens LFR component safety. BCC metals exhibit high LME sensitivity, while FCC metals show lower susceptibility. LME in LBE is characterized by reduced ductility and quasi-cleavage fracture morphology. [FIGURE:11] shows typical quasi-cleavage features on T91 fracture surfaces after LCF and fatigue crack propagation tests at 350°C in saturated-oxygen LBE [46,59].

Various models have been proposed for solid/liquid metal systems: adsorption-induced surface energy reduction, adsorption-induced atomic bond weakening, adsorption-enhanced dislocation emission, adsorption-promoted work hardening, and stress-assisted dissolution [2]. T91 exhibits severe embrittlement in specific LBE environments, with macroscopic transgranular quasi-cleavage and microscopic cracking along deformation-induced boundaries and high-dislocation-density interfaces. No single model fully explains T91 LME.

Gong et al. [45,51] proposed that LBE adsorption weakens atomic bonds at crack tips while dislocation pile-up causes hardening; when the critical stress for dislocation motion exceeds the weakened bond strength, cleavage occurs. Xue et al. [46,59] demonstrated that fatigue cracks in T91 at 150–450°C are macroscopically transgranular but microscopically follow 1–5° subgrain boundaries in the crack-tip plastic zone. Atomic-scale characterization [59] revealed LBE wetting of the plastic zone, with Pb-Bi segregation, short-range ordered Pb/Bi-Fe superstructures, and Pb-Bi clusters/precipitates/films along subgrain boundaries, reducing cohesive strength and initiating microcracks. These atomic-scale observations elucidate the quasi-cleavage LME mechanism. At 550°C in saturated-oxygen LBE, dynamic recrystallization at fatigue crack tips causes severe intergranular oxidation with Pb-Bi penetration, leading to oxidation+penetration-dominated LME [48].

Although 316L(N) SSRT and LCF tests show ductile fracture features with low LME sensitivity, Xue et al. [90] reported increased fatigue crack propagation rates and quasi-cleavage features in 316LN at 300–400°C in saturated-oxygen LBE. This likely results from crevice effects reducing crack-tip oxygen concentration [91] and plastic zone formation facilitating LBE wetting and stress-assisted penetration [92,93]. [FIGURE:13] shows Pb-Bi preferential diffusion along deformation twin boundaries and dislocation interfaces with Ni-selective dissolution in 316LN at 400°C, forming Pb₇Bi₃ particles, segregation, microvoids, and interfacial cracking that reduce cohesive strength [90]. Serre et al. [83] also observed reduced elongation and intergranular LME in AFA steels above 500°C. Luo et al. [94] demonstrated Bi segregation at Ni grain boundaries forming bilayer structures that embrittle Ni. Yu et al. [95] observed segregation-induced ordered superstructures at Ni-Bi grain boundaries causing embrittlement. Duscher et al. [96] showed Bi segregation at Cu grain boundaries induces electronic structure changes resembling Zn, causing embrittlement. Thus, despite low LME sensitivity in FCC alloys, Pb-Bi penetration and segregation along grain boundaries, twin boundaries, and dislocation interfaces can still cause quasi-cleavage or intergranular fracture.

5. Summary and Outlook

1. 9–12Cr ferritic/martensitic steels and 316 stainless steel are preferred materials for LFR fuel cladding and vessels. However, Fe₃O₄ and FeCr₂O₄ scales formed at high temperatures (>450°C) provide insufficient protection, leading to severe oxidation with intergranular attack under high oxygen and dissolution corrosion under low oxygen, further exacerbated by high flow velocities (~2 m/s).

2. Adding Si or Al to 9–12Cr ferritic/martensitic and austenitic stainless steels promotes formation of Si-rich protective scales, leading to development of Si-enhanced ferritic/martensitic steels (1–2 wt.% Si), Si-enhanced austenitic stainless steels (2.5–5 wt.% Si), and AFA steels (2–6 wt.% Al). However, dissolution corrosion risks remain under high-temperature, low-oxygen (<10⁻⁸ wt.%) conditions.

3. 9–12Cr ferritic/martensitic steels exhibit high LME sensitivity, typically showing quasi-cleavage (≤450°C) or intergranular (~550°C) fracture under load, degrading ductility, fracture toughness, fatigue, creep, and crack propagation. The quasi-cleavage mechanism involves LBE wetting of the crack-tip plastic zone, forming Pb-Bi segregation, ordered superstructures, and clusters along subgrain boundaries that reduce cohesive strength. High-temperature intergranular fracture results from combined intergranular oxidation and LBE penetration.

4. 316L(N) stainless steel shows low LME sensitivity with ductile fracture under load, maintaining ductility, fracture toughness, and fatigue performance comparable to air, though quasi-cleavage occurs during fatigue crack propagation via Pb-Bi penetration and segregation along grain boundaries, twin boundaries, and dislocation interfaces.

5. Current LME mechanisms are based on post-test microstructural analysis. Future work should employ in-situ TEM testing of LBE-wetted materials combined with molecular dynamics and first-principles calculations to investigate Pb-Bi interactions with vacancies, dislocations, and grain boundaries under load, fundamentally revealing LME mechanisms.

6. Limited data exist on creep, fatigue, and creep-fatigue of ferritic/martensitic and austenitic steels at typical service temperatures (500–550°C) in LBE. Intergranular oxidation at high temperatures may degrade performance through corrosion-stress interactions, requiring systematic investigation.

7. Si and Al are ferrite stabilizers that can reduce fracture toughness, increase ductile-to-brittle transition temperature, and degrade irradiation embrittlement resistance. While corrosion behavior of Si-enhanced and AFA steels has been studied, their performance under simulated irradiation-thermal-mechanical-corrosion coupling requires systematic evaluation to support engineering applications in LFRs.

Author Contributions: Jibo Tan conceived the overall idea and wrote the manuscript; Xinrui Zhang contributed the corrosion section; Baoquan Xue contributed the embrittlement mechanism section; Ziyu Zhang reviewed and edited the manuscript; Xinqiang Wu supervised the entire work and performed final review and editing.

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