Morphology and structure of the oxide layer

The oxide layer formed in the primary water of the simulated PWR had a double-layered structure. The inner layer was a continuous and dense Cr-rich oxide, and the outer layer was either an Fe-rich oxide (in austenitic stainless steel) or a Ni-rich oxide (in a Ni-based alloy). Both layers have spinel structures [34–41]. The thickness of the oxide layer varied from 0.1 to several μm. Figure 1 shows a schematic of the oxide layer formed on the surfaces of the stainless steel and Ni-based alloys in the simulated PWR. For austenitic stainless steel, as shown in Table 4, the outer oxide layer was composed of magnetite Fe₃O₄ or iron-nickel spinel oxide, whereas the chemical and crystal structures of the inner oxide layer remain controversial. The inner oxide layer was composed of either a Cr-rich face-centered cubic (fcc) phase [36,38,40,42], a hexagonal Cr₂O₃ [43-45], or a two-phase mixture [41,46]. Several studies have observed an enrichment of Ni at the oxide/alloy interface due to the selective oxidation and preferential movement of Cr and Fe from the alloy [34,35,38,47]. The Cr content of the alloy determined the thickness of the oxide layer [35], as shown in Fig. 2. Generally, the higher the Cr content of the alloy, the thinner was the oxide layer that was formed [35]. For Ni-based alloys, the outer oxide layer consisted of dispersed nickel ferrite (Ni_(1−y)Fe_(2+y)O₄) and nickel hydroxide. The metastable solid nickel hydroxide and ferric hydroxide formed on the surface could be caused by the precipitation of stable neutral aqueous complexes[48–51]. The reaction can be expressed as follows [48–51]: $\text{Ni}{\left(\text{OH}\right)}_{2\left(\text{aq}\right)}\rightleftarrows \text{Ni}{\left(\text{OH}\right)}_2$ (1) $\text{Fe}{\left(\text{OH}\right)}_{2\left(\text{aq}\right)}\rightleftarrows \text{Fe}{\left(\text{OH}\right)}_2$ (2) $\left(2+y\right)\text{Fe}{\left(\text{OH}\right)}_2+\left(1-y\right)\text{Ni}{\left(\text{OH}\right)}_2\overrightarrow{\leftarrow }{\text{Ni}}_{\left(1-y\right)}{\text{Fe}}_{\left(2+y\right)}{\text{O}}_4+2{\text{H}}_2\text{O}+{\text{H}}_{2\left(\text{aq}\right)}$ (3)

Fig. 1Full-screen View

(Color online) Schematic of the oxide layer formed on the surface of (a) stainless steel and (b) Ni-based alloys [50] in a simulated PWR environment. x and y represent coefficients that are related to the number of Fe, Ni, and Cr atoms in the studied medium

Fig. 2Full-screen View

Effect of Cr content on the inner oxide film thickness exposed in the simulated PWR primary water at 320 ℃ for 380 h [35]

The inner oxide layer was composed of Fe-rich and contained nickel chromite (Ni_(1−x)Fe_xCr₂O₄). As can be seen from Fig. 1(b), Cr₂O₃ particles were unevenly distributed at the interface between the inner oxide film and the substrate. The number of Cr₂O₃ particles was related to the defects on the substrate surface. Some studies have also observed that there is no chromium depletion at the alloy/oxide interface [52–54]. Nevertheless, a Cr-depleted zone was observed beneath the oxide scale. A disturbed microstructure of the alloy with small grains and a strong dislocation density was observed in this Cr-depleted layer [52,54].

Formation and growth of the oxide layer

The passivation process of metals in a solution is initiated with the adsorption of anions, especially OH^–. Previous studies [55-57] have indicated that the process of forming an oxide layer on the surface of the Fe-Cr-Ni alloys, such as stainless steel, in a simulated PWR mainly relies on the diffusion of the alloying elements Fe, Cr, and Ni. The growth of the outer oxide layer formed on the alloy surface occurs via the dissolution and reprecipitation mechanism, whereas that of the inner oxide layer occurs via the solid growth mechanism [39,42,55]. Figure 3 schematically illustrates the oxide formation process on the stainless steel exposed to the environment inside a simulated PWR. During the initial stage of corrosion, when the material is exposed to high-temperature and high-pressure water, the oxygen in the environment comes into direct contact with the metal surface of the material. Among the alloying elements, Cr exhibits the strongest affinity for oxygen to form oxides. Consequently, during the initial stage, the material surface gradually develops a Cr-enriched oxide film. Due to the diffusion rate in oxides, Fe > Ni >> Cr [39,42,55], the comparatively faster diffusion rate of Fe results in its diffusion outward into the aqueous medium or its interaction with adsorbed H₂O, eventually precipitating into the outer oxide particles. At the same time, external oxygen diffuses into the interior of the matrix material and reacts with Cr to form an inner oxide film. Considering that the diffusion rate of Ni in the inner oxide film lies between that of Cr and Fe, Ni neither accumulates extensively within the inner layer nor exists in significant quantities within the outer oxide layer. Consequently, the outer oxide layer is predominantly composed of Fe, whereas the inner oxide film is enriched with Cr. Thin Ni enrichment may occur at the metal oxide/interface, depending on the diffusion rate of Ni.

Fig. 3Full-screen View

(Color online) Schematics showing the formation process of the oxide on stainless steel exposed to a simulated PWR environment

The factors affecting the growth of the oxide layers in the simulated PWR are summarized in Table 5. The chemical parameters of water (DH [29,58-63], DO [29, 64-66], and PH[56, 67,68]) are the key factors affecting the corrosion resistance and properties of the oxide in a PWR primary system. Among these parameters, the DH is a favorable factor for reducing the oxidation rate. In particular, the primary water undergoes radiolysis when it is subjected to irradiation, producing oxidizing species (O₂ + H₂O₂) and making the corrosion electric displacement more positive, which, in turn, increases the corrosion rate. Irradiation affects the corrosion behavior of materials by changing the water environment, as will be discussed in detail in Sect. 2.4.3.

Matthews et al. [69] suggested that the reaction rate at the interface is the principal factor affecting the oxide growth, surpassing other contributing factors. The oxide growth depends on the interfacial reaction rate when the same sample is simultaneously exposed to the PWR environment for the same time. According to the point-defect model proposed by Macdonald et al. [71], in high-temperature and high-pressure water environments, oxygen ion vacancies can be formed at the interface between the inner oxide layer and the base metal through reaction (4) [51,71]. The metal atoms are converted into metal-ions via reactions (4)–(6) [51,71,74]. $m\leftrightarrow m_{\text{m}\left(\text{oxide}\right)}^x+V_{\text{O}\left(\text{oxide}\right)}^{\mathrm{..}}+2{\text{e}}^{-}$ (4) $\text{Ni}\leftrightarrow {\text{Ni}}^{\mathrm{..}}+2{\text{e}}^{-}$ (5) $\text{Fe}\leftrightarrow {\text{Fe}}^{\mathrm{..}}+2{\text{e}}^{-}$ (6) where m denotes the metal atom in the base metal, namely Cr, Ni, and Fe, $m_{\text{m}\left(\text{oxide}\right)}^x$ are normal metal-ion lattice sites in the inner oxide, $V_{\text{O}\left(\text{oxide}\right)}^{\mathrm{..}}$ represents the oxygen-ion vacancy in the oxide, and ${\text{Ni}}^{\mathrm{..}}$ and ${\text{ Fe}}^{\mathrm{..}}$ are the interstitial metal cations in the inner oxide layer.

O^2- diffuses to the oxide/matrix interface (oxygen-ion vacancies outward surface) and combines with metal cations to form the inner oxide layer, resulting in the inward migration of the oxide/matrix interface. The oxygen vacancy and cation diffusion rates control the growth rate of the inner oxide. Further, the metal cations formed in the inner oxide layer diffuse into the aqueous medium or interact with the absorbed H₂O and eventually precipitate into oxide particles driven by the dissolution and redeposition mechanisms [40,65]. The diffusion rate of cations in the inner oxide layer plays a major role in the growth of the outer oxide layer. Irradiation controls the growth of the oxide by affecting the diffusion rate, as described in Sect. 2.4.

Effect of irradiation defects on corrosion

After irradiation, defects such as voids, helium bubbles, black spots, and dislocation rings are produced inside the material [75,81]. For a comprehensive understanding of the composition and structure of the oxide formed on the metal surface, as well as the depth of the inner oxide layer following irradiation, extensive studies have been conducted, which have been summarized in Table 6[19–21,76,82,83]. Although Deng et al.[20] have already reported that a finer oxide grain in the inner oxide film was formed on a 3-dpa proton-irradiated 316 L austenitic stainless steel after exposure to the primary water of the simulated PWR for 500 h, it is generally accepted that irradiation does not alter the crystal structure and the qualitative chemistry of the inner and outer oxides formed on materials, which can be characterized using high-resolution TEM (HRTEM) [19,21,71, 77,78, 82,84,85]. Similar to the unirradiated sample, the oxide layer maintains its characteristic double-layer structure when stainless steel is irradiated. The outer layer comprises polyhedral oxide particles primarily composed of Fe-rich oxides, such as dispersed Fe-Ni spinel or magnetite Fe₃O₄, and the inner continuous and protective oxide film is predominantly composed of Cr-rich oxide, known as Fe-Cr-Ni spinel. Both layers have been characterized as spinel structures. In addition, Ni enrichment at the oxide/matrix interface has been observed [19,82]. However, based on the available study data [19-21,74,76,82,85-89], it has been observed that irradiation affects the quantitative chemistry, morphology, and thickness of the oxide layers formed on the alloys. In irradiated materials, the Cr content in the inner layer has been observed to be high [19, 76-78,85]. In contrast, a lower Cr content has been observed in the inner oxides of 316 L stainless steel irradiated with protons exposed to high-temperature water, due to the dissolution of Cr-rich spinel oxides in irradiated water with added hydrogen [90]. Further, no significant change in the Cr concentration in the inner oxide has been reported with increasing irradiation dose [20,21,74]. These different oxidation behaviors probably arise from the different irradiation and oxidation conditions (dose, temperature, oxidation duration, etc.) In addition to the change in the Cr concentration, Kuang et al. [78] observed that the maximum Ni content in the transition zone is higher and wider.

Generally, these irradiation-induced defects accelerate corrosion primarily in two ways [19,20,74,83,85-89]: (i) by increasing the size and/or density of the outer oxide particles and (ii) by increasing the depth of the inner oxide layer. The oxidation kinetics model has been employed to elucidate the changes in the inner oxide layer thickness caused by irradiation-induced defects [74,89]. Based on the quasi-stationarity assumption [89], the thickness of the oxide scale on the surface of a nickel-based alloy in a PWR can be described using Eq. (7) as follows: $X\left(t\right)=\sqrt{\frac{\delta \times D_{\text{sc}}^{\text{O}}}{\phi }\times t+{\left(X_0\right)}^2}$ (7) where X(t) represents the thickness of the inner oxide layer, $X_0$ denotes the thickness of the initial oxide film, t is the oxidation time, $\delta \text{ and }\phi$ are the width and size, respectively, of the grain boundaries in the oxide. $D_{\text{sc}}^{\text{O}}$ is the diffusion coefficient of oxygen vacancies along short circuits (such as the grain boundaries in oxides). Only the grain boundaries act as channels for the diffusion of oxygen vacancies in the absence of defects. However, the grain boundaries as well as the irradiation-induced defects contribute to the diffusion of oxygen vacancies when subjected to irradiation. Therefore, it is anticipated that $D_{\text{sc}}^{\text{O}}$ will increase, indicating a higher diffusion rate of oxygen vacancies, resulting in the formation of a thicker inner oxide layer, according to Eq. (7).

On the other hand, the irradiation-induced defects are the preferred nucleation points for the oxide particles, and thus the number of outer oxide particles is higher. These defects function as fast diffusion channels, facilitating the outward diffusion of metal cations. Thus, the oxide particles formed on the surface of the material are larger than those in the unirradiated state. The oxidation kinetics provide insights into the oxidation rate. Higher oxidation kinetic values in the irradiated region indicate that irradiation accelerates the corrosion process. For example, Boisson et al. [19] reported a five-fold increase in the oxidation kinetics of irradiated 316L austenitic stainless steel compared to the unirradiated area after 24 h of oxidation in a simulated PWR. Similarly, Deng [20] emphasized that radiation enhances the oxidation kinetics (up to four times faster) during a longer oxidation duration about 500 h in irradiated 304 nuclear-grade stainless steel. Thicker inner oxide layers in the irradiated areas indicate faster oxygen diffusion. However, Perrin et al. [76] conducted an experiment where they irradiated 316L stainless steel with protons and subjected it to corrosion in a simulated PWR at 325 ℃ for 1024 h. Surprisingly, they found that oxygen diffusion in the inner oxide layer of the irradiated sample occurred at a slow pace. The diffusion coefficients in the irradiated sample and unirradiated sample were 5×10^-17 cm²s^-1 and 3×10^-16 cm²s^-1, respectively. The outer oxide particles formed on the irradiated sample were smaller in size, and the inner oxide film was thinner than that of the unirradiated sample, as shown in Fig. 4. These observations indicate that irradiation inhibits corrosion. The growth of the inner oxide film was constrained by the slow diffusion of oxygen through the inner oxide layer. Irradiation-induced defects promoted the diffusion of Cr, resulting in the formation of a thicker Cr-rich film. A dense Cr-rich film normally acts as a protective layer that impedes oxygen diffusion and further oxidization. Kuang et al. [78] also suggested that irradiation-induced defects enhance the diffusion of Cr, resulting in higher protection of the inner oxide layer, which is one of the reasons for the enhanced oxidation resistance in the irradiated region. Further, they observed that the continuous Ni-rich zone formed near the oxide/matrix interface reduces the "available space" for oxidation. It should be noted that in the experiment conducted by Perrin et al., the samples were only mechanically polished. Samples in which cold work is induced by polishing [91,92], tend to exhibit faster oxidation kinetics. However, this cold working diminishes under irradiation [93] due to the disappearance of the initial dislocation network. Consequently, the irradiated area without cold work experiences a slower rate of oxygen diffusion, resulting in thinner oxides. Therefore, sample polishing must be strictly controlled, and the results obtained by Perrin et al. by irradiating samples with protons (or neutrons) using mechanical polishing need to be validated to determine whether polish-induced cold working overcomes the effects of irradiation.

Fig. 4Full-screen View

Observation of the oxide in the (a), (c) unirradiated and (b), (d) irradiated 316L stainless steel corroded for 1040 h at 325 ℃ in the primary water of the simulated PWR [76]

The corrosion of the alloy intensifies with increasing corrosion time, as shown in Fig. 5(a) [82,89]. Furthermore, the density and/or size of the irradiation-induced defects are influenced by the irradiation dose or temperature. For a given material, the irradiation-induced defect density and/or size increases with increasing irradiation dose (Fig. 5(b)) or elevated irradiation temperature [20,74,86]. Consequently, the diffusion rates of the metal cations and oxygen anions increase, resulting in an elevated oxidation rate. It is worth noting that a smoother surface appears in the irradiated alloy due to irradiation-induced ion sputtering [85], as shown in Fig. 6(c). During the corrosion process, rough surfaces with large surface areas absorb more dissolved oxygen and impurities, thereby enhancing the nucleation rate of oxides. In contrast, the irradiated alloys exhibit fewer oxide particles. Simultaneously, high-energy particle irradiation causes localized temperature increases on the material surface and the formation of Cr-rich oxides (Fig. 6 (d)). The thin layer composed of Cr-rich oxide prevents the interaction between the underlying elements and external oxygen atoms, thus hindering the growth and nucleation of iron-rich oxides on the alloy surface in the initial corrosion stage (Fig. 6 (a)). In addition, the short-term diffusion of atoms becomes difficult due to defects. However, for long-term corrosion, irradiation-induced defects promote the long-range migration of metal atoms, and thus the oxidation rate and corrosion weight gain of the irradiated alloys are higher than those of the irradiated alloys (Fig. 6(a) and 6(b)). The defects are principally point defects. Their impact on the diffusion of enhanced oxygen and metal atoms during the initial stage of corrosion is minimal, and the corrosion behavior of materials is almost unaffected by irradiation [77,79]. Jiao et al. [21] observed that proton-irradiated SUS316 stainless steel corroded for 70 h and that the thickness of the inner oxide film was only marginally affected by irradiation, as shown in Fig. 7. However, no conclusive explanation for this phenomenon was given. Further empirical research is necessary to investigate the influence of irradiation-induced defects on material corrosion, including exploration of the critical irradiation dose that either promotes, inhibits, or has no effect on material corrosion, as well as the influence and mechanism of irradiation on material corrosion at the initial stage of irradiation. This will enable obtaining a more systematic understanding of the effect of irradiation on corrosion, which has guiding significance for the application of nuclear materials in the nuclear industry.

Fig. 5Full-screen View

(Color online) Oxide thickness as a function of the (a) corrosion time [82,90] and (b) irradiation dose [20,71,87] in a simulated PWR environment

Fig. 6Full-screen View

(Color online) (a) Histogram of weight gain of the Fe₁₆Cr_4.5Al alloy samples with corrosion time. (b) SEM images showing the surface morphologies of the Fe₁₆Cr_4.5Al alloy samples corroded for 50 h and 200 h (b1, b2) as-received and irradiated to (b3, b4) 60 dpa. (c) SEM images showing the surface morphologies of the Fe₁₆Cr_4.5Al alloy samples: (c1) as-received sample and (c2) magnified images of the rectangular region in (c1), and samples irradiated to (c3) 10 dpa and (c4) 60 dpa. (d) Cross-sectional TEM bright-field images of the Fe₁₆Cr_4.5Al alloy samples after (d1) and (d2) irradiation to 60 dpa [85].

Fig. 7Full-screen View

(a) Surface morphology of the oxide scale formed on SUS316 stainless steel with irradiated and unirradiated areas. TEM observation of the (b) unirradiated and (c) irradiated samples [21]

Effect of radiation-induced segregation on corrosion

Radiation-induced segregation (RIS) refers to the phenomenon of microchemical segregation at the grain boundaries (GB) due to irradiation, leading to the depletion of Cr, Fe, and Mo, and the enrichment of Si, Ni, and P [15,94,95]. Post-irradiation annealing eliminates element segregation at the grain boundary so that no local corrosion of the grain boundaries is observed [88]. This indicates that RIS is primarily responsible for the intergranular oxidation of materials. Grain boundaries act as the preferred pathways for diffusion and are therefore more susceptible to oxidation. Intergranular corrosion occurs in the irradiated as well as unirradiated areas, but the deeper oxide penetration in the irradiated area suggests that irradiation intensifies intergranular corrosion [19,20,74,86,96-98], which is attributed to the depletion of Cr at the grain boundaries. The lower Cr content in the grain boundary results in the increased penetration of the oxidation depth (Figs. 8 and 9) and decreases the intergranular corrosion resistance [20,86]. Furthermore, the unirradiated alloy exhibits Ni-rich regions before the oxide tip at the grain boundary, whereas such regions are not formed in the irradiated samples [20]. This observation further supports the notion that irradiation promotes a high intergranular corrosion rate. The selective oxidation of Cr and Fe leads to the preferential removal of these elements from the alloy, forming Ni-rich regions. In irradiated alloys, rapid corrosion along the grain boundaries does not provide sufficient time for selective oxidation to occur [20,99].

Fig. 8Full-screen View

(Color online) Depth of the intergranular oxide and Cr segregation by RIS as a function of the irradiation dose [20,87]

Fig. 9Full-screen View

(Color online) Cr segregation as a function of the irradiation dose [100–102]

The depth of the intergranular oxidation increases with increasing irradiation dose (Fig. 8) [20,103], possibly because of maximum Cr consumption at the grain boundaries at higher doses. Studies [100–102] have shown that the depletion value of Cr at the grain boundary changes dramatically at low and medium doses (0.1~5 dpa) but no longer changes significantly with further increase in the radiation dose, or even remains unchanged at higher doses, as shown in Fig. 9. Fukumura et al. [86] found that the degree of Cr segregation in neutron-irradiated stainless steel exposed to a simulated PWR for 1149 h did not change significantly with the increase in dose from 25 dpa to 73 dpa, but the oxidation length at the grain boundary increased by approximately 20%, as depicted in Fig. 8(b). It is challenging to quantitatively explain the promotion of oxidation at the grain boundary at high irradiation doses through Cr depletion alone. The evolution of Ni and Si segregation at the grain boundaries may be implicated in the promotion of oxidation at these boundaries at high irradiation doses (Fig. 8(b)).

However, the presence of Si segregated at the irradiated grain boundary appears to inhibit intergranular corrosion. Kuang et al. [94] compared the intergranular oxidation behavior of solution annealing (SA) and 2.5 dpa proton-irradiated specimens of 316L in a simulated PWR with exposure 1000 h, and found that the penetration depth of the intergranular oxide in the irradiated area was shallower than that in the unirradiated area. This suggests that irradiation enhances the intergranular oxidation resistance of stainless steel. Figure 10 shows the intergranular oxidation process at the irradiated and unirradiated grain boundaries [94]. Researchers have proposed that Si separated at the grain boundaries of the irradiated samples tends to diffuse outward and undergoes oxidation due to its high diffusion coefficient and affinity for oxygen. The Si-rich oxide at the intergranular oxide tip of the irradiated sample serves as a temporary diffusion barrier for oxygen, although it tends to dissolve near the sample surface. The migration efficiency of the elements (especially Cr) along the grain boundary of the irradiated sample to the oxidation front improves. This is because the preferential diffusion of Si generates vacancies, which can promote the diffusion of Cr atoms and lead to an increase in the Cr content at the tips of the intergranular oxides. Irradiation-induced structural defects also contribute to the diffusion of elements such as Cr along the grain boundary. Therefore, although Cr in the grain boundaries of the irradiated samples is depleted due to RIS at the original grain boundaries, Cr is enriched in the oxides along the grain boundaries of the irradiated samples. The coenrichment of Si and Cr at the tips of the intergranular oxides enhances oxidation resistance, ultimately leading to a decrease in the oxidation rate of the irradiated grain boundary. In addition, the grain boundary structure affects RIS [104,105] and intergranular corrosion [106]. Therefore, future experimental analyses should focus on selecting the same type of grain boundary across irradiated and unirradiated areas. This will enable direct comparative studies of the intergranular oxidation behavior between the two areas at a single grain boundary, eliminating the interference caused by the differences in the GB structure differences. This is crucial to for distinguishing the effects of irradiation on intergranular oxidation.

Fig. 10Full-screen View

(Color online) Schematic showing the oxidation behavior at the grain boundaries of (a) irradiated and (b) unirradiated 316L exposed for 1000 h in the primary water environment of a simulated PWR [94]

In summary, irradiation segregation leads to Cr depletion and Ni and Si enrichment at the grain boundaries. However, its effect on the intergranular oxidation of materials remains unclear. In addition, the impact of RIS on intergranular corrosion has mostly been investigated at low radiation doses, but there is a lack of research at high radiation doses. Future studies should aim to strengthen our understanding of the mechanism of intergranular corrosion of materials at high radiation doses to bridge this research gap.

Effect of irradiation radiolysis on corrosion

Irradiation can directly affect material corrosion through irradiation damage, as described in detail in Sects. 2.1 and 2.2. Furthermore, radiolysis can affect corrosive media. In PWRs, neutrons interact with the water coolant, leading to radiolysis (Eqs. (8) and (9)). This process results in the formation of long-lived molecules (H₂, O₂, and H₂O₂) and short-lived free radical substances (${\text{H}}^{.}$, ${\text{OH}}^{.}$, ${\text{e}}_{\text{aq}}^{-}$, and ${\text{HO}}_2^{\cdot }$). These species can undergo oxidation (O₂，H₂O₂，HO₂) or reduction (${\text{e}}_{\text{aq}}^{-}$, H, and H₂). Literature [100,107–111] provides a detailed review of the important reactions associated with these species and explains the charge conservation of H⁺ and OH^-. ${\text{H}}_2\text{O}\stackrel{\text{Ionizing adiation}}{\to }{\text{e}}_{\text{aq}}^{-},{\text{H}}^{\cdot },{\text{HO}}^{\cdot },{\text{HO}}_2^{\cdot },{\text{H}}_3{\text{O}}^{+},{\text{OH}}^{-},{\text{H}}_2,{\text{H}}_2{\text{O}}_2$ (8) ${\text{H}}_2{\text{O}}_2\to {\text{O}}_2+2{\text{H}}^{+}+2{\text{e}}^{-}$ (9)

The water radiolysis product H₂O₂ affects the electrochemical corrosion potential [108,110,111]. Figure 11 [112] shows the change in the corrosion potential of 304 and 316 stainless steels as a function of the H₂O₂ concentration. As the H₂O₂ concentration increases, the corrosion potential also increases until it reaches a saturation point. The saturated corrosion potential can extend to the penetration zone of the passivated film, where cracking/dissolution of the oxide film occurs [113-114]. When subjected to irradiation, the corrosion potential increases significantly. Was et al. [115,116] performed thermodynamic calculations and radiolysis modeling and found that at very high irradiation dose rates (4000 kGy/s), the corrosion potential increased by ~0.8V_SHE compared to the unirradiated samples, thereby promoting material corrosion. In addition, in high-temperature water with low dissolved oxygen concentration and no dissolved hydrogen, the corrosion potential increased by more than 0.25V_SHE.

Fig. 11Full-screen View

(Color online) Corrosion potential curve of 304 and 316 stainless steel as a function of the H₂O₂ concentration [112]

The characteristics of the oxides formed on the surface of the irradiated materials provide insight into the alteration of the stability of the species within the oxide film due to the irradiated environment. In experiments, the materials were irradiated with a proton beam and simultaneously exposed to a simulated PWR. Although radiolysis and irradiation damage occur on the surface of the materials simultaneously, in a series of experiments conducted by Was et al. [90,115,117], stainless steel samples not directly irradiated but exposed to radiolysis water exhibited oxide characteristics similar to those of the irradiated samples. Hematite was observed on the surface of the oxide, and Cr was deficient in the inner oxide layer, which differed from the oxide characteristics of the unirradiated sample. This indicated that water radiolysis was the primary mechanism affecting the properties and formation of the oxide. Comparing the oxides in the irradiated and unirradiated areas, such as the effect of irradiation-induced defects on corrosion, irradiation did not alter the double-layer structure of oxides formed on the material surface, which still consisted of an outer layer of oxide particles and an inner layer of a continuous oxide film. In situ irradiation corrosion experiments demonstrated significant changes in the morphology, composition, and depth of the inner oxide film compared to the unirradiated samples [87,115–119]. Using austenitic stainless steel as an example, the following observations were made:

•　The oxides formed on the surfaces of the irradiated samples predominantly consisted of hematite and relatively few large planar oxide particles, primarily consisting of densely packed small equiaxed crystals.

•　The irradiated samples displayed a thinner and more porous inner oxide film with a lower Cr content, indicating the loss of Cr from the inner oxide layer after irradiation.

•　The interface between the outer particles and the inner oxide was indistinct, and the metal/oxide interface fluctuated significantly in the irradiated samples.

As mentioned earlier, the long-lived radiolysis products (H₂O₂) generated by irradiation can significantly increase the corrosion potential. This alteration in the potential leads to the oxidation of Fe²⁺ to Fe³⁺ and the subsequent formation of hematite. Cr-rich spinel oxide dissolves to form ${\text{HCrO}}_4^{-}$ due to the increased corrosion potential. Consequently, the thickness of the inner oxide layer decreases, and Cr loss occurs. A lower Cr content indicates a less protective oxide film, leading to porosity. A higher diffusion rate of oxygen vacancies and cations due to worse protection by the inner scale results in further corrosion, contradicting the theory that the inner oxide layer is thinner. Hence, the growth of the inner oxide film on the irradiated alloy at a high-temperature and high-pressure water can be characterized as a competing process of solid oxidation and dissolution of the existing film.

Thermodynamic driving corrosion

The coolant salts used in an MSR typically consist of mixtures of alkali metal or alkaline earth metal fluorides such as LiF-NaF-KF or LiF-BeF₂. The driving force for molten salt corrosion originates from the difference in the Gibbs free energy between the molten fluoride salt and the metal fluoride. It is widely known that alkali metal (or alkaline earth metal) fluorides have the lowest Gibbs free energies among metal fluorides. Consequently, alloying elements, such as Cr, Fe, and Ni, cannot be chemically replaced by reactions in alkali metal (or alkaline earth metal) fluorides. Further, thermodynamic calculations indicate that the equilibrium concentrations of metals such as Ni, Mo, Co, Fe, and Cr in corrosion reactions involving LiF, NaF, KF, and BeF₂ are extremely low [144]. In general, when the equilibrium concentration of the reaction products is very low, it can be considered that the reaction does not occur. Therefore, the intrinsic corrosion of alloys containing these major elements in the impurity-free molten salt coolants is almost impossible. First-principles simulations and Molecular dynamics simulations have been used to investigate the thermodynamic properties of the molten salt. The results thus obtained reveal that the molten salt composition plays an important role in the corrosion of Ni-based alloys in molten fluoride salts [145-147]. These studies help us to efficiently understand the basic fluorine-induced initial corrosion mechanism of Ni-based alloys in molten salt environments.

Impurity-driven corrosion

The impurities present in the molten salt are the main inducers of corrosion, especially in the initial stages, as shown in Fig.12. Oxidizing impurities (H₂O, HF, NiF₂, FeF₂, metal oxides, oxyacid ions, etc.) can act as oxidants in molten salts, leading to the rapid corrosion of alloys through various reaction mechanisms (Eqs. (10)–(14)) [120, 124]. However, as these impurities are consumed, the corrosion rate decreases, transitioning to a slower stage of corrosion. ${\text{ H}}_2\text{O}\left(\text{g}\right)+2\text{MeF}={\text{M}}_2\text{O}+2\text{HF}\left(\text{g}\right)$ (10) $y\text{HF}\left(\text{g}\right)+x\text{Me}={\text{Me}}_x{\text{F}}_y+\frac{1}{2}y{\text{H}}_2\left(\text{g}\right)\left(\text{Me=Cr, Fe, Ni, Co, Al}\right)$ (11) $\begin{array}{ll}{\text{FeF}}_2+\text{Cr}={\text{CrF}}_2+\text{Fe}\hfill & 2{\text{FeF}}_3+3\text{Cr}=3{\text{CrF}}_2+2\text{Fe}\hfill \end{array}$ (12) $\begin{array}{ll}{\text{NiF}}_2+\text{Cr}={\text{CrF}}_2+\text{Ni}\hfill & 2{\text{CrF}}_2+{\text{NiF}}_2=2{\text{CrF}}_3+\text{Ni}\hfill \end{array}$ (13) ${\text{SO}}_4^{2-}+2\text{Cr}=\text{S}+{\text{O}}^{2-}+{\text{Cr}}_2{\text{O}}_3$ (14)

Fig. 12Full-screen View

(Color online) Driving force of molten salt corrosion at different corrosion stages

In the complex environment of nuclear reactors, the corrosion of alloys in molten salt environments is affected by multiple factors, including temperature differences [148,149], variations in the molten salt flow rate [150-152], presence of fission products [140-142], heterogeneous material corrosion (galvanic corrosion) [124,143,153-155], and irradiation-induced corrosion. In this paper, we focus on the effect of irradiation on the corrosion of alloys in molten salt environments and discuss the latest research progress.

Helium bubbles

A critical problem for Ni-based alloys in an MSR is the accumulation of helium atoms [164,165]. The ⁵⁹Ni nuclide in the Ni-based alloy has a very high (n, α) reaction cross-section, leading to the generation of a large amount of helium. Another important source of helium generating is the (n, α) reaction of thermal neutrons with ¹⁰B impurities in the alloy in the early stage of nuclear reactor operation. These He atoms can easily gather around the grain boundary to form He bubbles or cavities and cause intergranular fracture. This reduces the toughness of the alloy and its long-lasting life significantly, and results in the He brittleness of the alloy [166]. Thus, helium irradiation experiments were conducted on UNS N10003 alloy in a molten salt environment [167-172]. The results showed that helium irradiation accelerated the corrosion of the UNS N10003 alloy at a certain dose, at which obvious helium bubbles were observed via TEM characterization [170]. The presence of helium bubbles accelerated the corrosion of the UNS N10003 alloy in a molten salt environment in the following ways: 1) accelerating the dissolution of the alloying elements [170,173,174], 2) coarsening the surface of the alloy and enhancing intergranular corrosion [167], and 3) accelerating the formation of holes on the surface of the alloy [169,170].

An important mechanism by which helium bubbles accelerate the molten salt corrosion of alloys is the migration and diffusion of helium bubbles. Zhu et al. observed that the size of helium bubbles increased from 0.8 nm to 20 nm after the alloy was immersed in a molten salt environment. The migration of helium bubbles toward the interior bulk of the sample was observed, which coarsened the holes on the alloy surface and enhanced the intergranular corrosion of the UNS N10003 alloy [167]. Liu et al. reported the growth and coalescence of helium bubbles around the peak damaged region, which had no significant effect on the loss depth of Cr [169]. Lei et al. observed a unique outward migration of helium bubbles in a helium-ion-irradiated UNS N10003 alloy in a high-temperature molten salt environment, as shown in Fig.13. Helium bubbles can modify surface morphology and form holes through their outward migration. The formation of holes on the surface increases the contact area between the molten salt and the alloy, thus increasing the dissolution of Cr [170]. The dissolution of alloying elements also creates vacancies for the growth of helium bubbles (Table 7).

Fig. 13Full-screen View

Outward migration and growth of He bubbles from inside to the alloy surface after molten salt corrosion

Displacement damage defects

Fast neutrons in reactors cause displacement damage to the alloy through cascading collisions and form irradiation-induced defect clusters inside the alloy. These defect clusters strongly impede dislocation motion, resulting in the reduction of uniform elongation and fracture toughness of the alloy and inducing its hardening and embrittlement. There is also evidence that displacement damage can influence the corrosion of alloys in molten salt environments.

Researchers from the Massachusetts Institute of Technology reported that irradiation can decelerate the intergranular corrosion of Ni-Cr alloys in molten salt via an in situ proton beam irradiation-molten salt corrosion experiment [175]. They attributed this inhibition to the self-healing of grain boundaries caused by the diffusion of Ni and Cr interstitial particles from proton irradiation. Irradiation introduces abundant interstitials into the alloys and enhances bulk diffusion. The diffusion of irradiation-induced interstitials to the grain boundary supplies dissolving atoms to the molten salt, thus decelerating the intergranular corrosion of the Ni-Cr alloys, as shown in Fig. 14. ORNL carried neutron irradiation- molten salt corrosion experiment on UNS N10003 alloy and 316SS alloy at 800 ℃ with an irradiation dose of 7.5×10¹⁶ n/cm^{2 [176]}. The composition of the molten salt was NaCl-MgCl₂. The results of this study showed that neutron irradiation did not significantly affect the corrosion of the UNS N10003 alloy but decelerated the corrosion of the 316SS alloy. Helium ion irradiation can introduce helium bubbles into alloys, which can cause displacement damage. Researchers from Shanghai Institute of Applied Physics (SINAP) investigated the effect of helium ion irradiation with low and high ion fluences on the corrosion behavior of Inconel 617 and GH3535 alloys in FLiNaK molten salt at 700 ℃ [172]. For the substrates of both alloys, corrosion was observed to be accelerated with increasing irradiation doses. For the grain boundaries of the 617 alloy, the inhibition and promotion of intergranular corrosion were observed at low and high irradiation doses, respectively, which could be ascribed to the blocking of interstitial atoms, promotion of molten salt diffusion by helium bubbles, and self-healing mechanisms, as shown in Fig.15.

Fig. 14Full-screen View

(Color online) Inhibition of corrosion of the NiCr alloy in FLiNaK salt following the in situ proton beam irradiation. (a) Corrosion depth of the alloy for different irradiation doses, and (b) schematic of the irradiation inhibition corrosion mechanism[175]

Fig. 15Full-screen View

(Color online) Synergistic mechanism of irradiation and corrosion of Inconel 617 alloy [172]

Current studies have investigated the synergistic effects of neutron irradiation and corrosion, in situ and off-site ion irradiation, and corrosion. However, the selected alloys, irradiation temperatures, irradiation doses, and molten salt systems are all different. More experiments still need to be performed to clarify the effect of displacement damage on the molten salt corrosion of alloys.

RIS

Particle irradiation can induce element segregation and local microstructural changes in materials and is usually defined as RIS. During irradiation, point defects (vacancies and interstitial atoms) diffuse into nearby traps, such as the surface or internal grain boundaries of materials. Generally, the migration rates of these point defects differ. Some migrate to the sinks, whereas others remain away from them. Therefore RIS causes significant changes in the local composition of the material and affects its macroscopic properties.

RIS and precipitation are common in the pressure vessel steel of PWR. At low doses, the segregation of Cu, which leads to precipitation, can be observed by small-angle neutron scattering and scanning TEM (STEM). In addition, the segregation of Ni, Mn, P, Si, and other elements at dislocations or grain boundaries has also been reported [177]. In Cr-containing austenitic stainless steel and nickel-based alloys, irradiation-induced segregation leads to the diffusion of Cr to the grain boundaries, thus causing IASCC of the material. At lower temperatures, the concentration of defects increases, but they migrate to the sinks, leading to point-defect recombination. At higher temperatures, thermal diffusion is the dominant mechanism, and the composition of the material tends to be uniform. In the intermediate temperature (0.3–0.6T_m), RIS is the most serious due to the “anti-Kirkendal effect.” When the irradiation dose reaches 1–10 dpa, RIS can lead to local corrosion, grain boundary embrittlement, and a decline in the mechanical properties of the material. When the irradiation dose reaches 10 dpa, RIS leads to phase instability in the material, causing the generation of a new phase or decomposition of the existing phase [178]. In the Ni-Si system, when the solubility of Si in Ni reaches a certain degree, a new Ni₃Si phase will be formed (γ’ Phase) [179]. An RIS phenomenon was also observed in the alloy used for the MSR. Liu et al. and others used 30 keV Ar ions to irradiate a UNS N10003 alloy at room temperature. When the dose reached 48.4 dpa, a high concentration of Ni was observed at the grain boundary [180]. Huang et al. found that the precipitates rich in Mo and Cr appeared in the UNS N10003 alloy when the dose reached 10 dpa after Xe ion irradiation at 650 ℃ [181].

In a helium ion irradiated and molten salt corroded UNS N10003 alloy, Zhu et al. observed significant segregation of Si at the helium bubble surfaces, as shown in Fig. 16. When the segregation of Si increased to values above the solubility limit of ~ 10 at. % at the bubble surfaces [168], precipitation of a Ni-Si intermetallic compound occurs. The segregation of Si promotes chemical reactions between the Si atoms and the molten salt, thus enhancing localized corrosion damage. However, the formation of Ni-Si precipitates at large bubbles results in galvanic corrosion due to the difference in the electrochemical potentials between the Ni matrix and Ni-Si precipitates. Thus, RIS accelerates the localized corrosion damage of Ni-based alloys exposed to high-temperature molten salts.

Fig. 16Full-screen View

(Color online) Distribution of chemical elements in the uncorroded region of a sample irradiated and then corroded for 200 h: (a) High-angle annular dark-field imaging, and STEM-Energy dispersive X-ray spectroscopy (EDS) maps of (b) Ni, (c) Si, (d) Cr, (e) Mo, and (f) Fe ^[168].

Effects of the irradiation environment

In an MSR, the presence of fission products can significantly affect the corrosion of structural materials. The accumulation of fission products in the molten salt, resulting from the fission reactions in the core region, can interfere with the corrosion process and potentially promote corrosion due to the diffusion of specific fission products into the alloy. One such example is the deposition and diffusion of the fission product tellurium (Te) into the UNS N10003 alloy, which can lead to embrittlement of the grain boundaries of the alloy and potentially induce stress corrosion cracking. The extent of grain boundary embrittlement cracking caused by Te depends on the diffusion rate of Te in the alloy. In addition, the occurrence of Te-induced stress corrosion cracking at the grain boundaries is influenced by the concentration of Te and the applied stress load. [182-186].

These observations highlight the influence of Te concentration and applied stress on the susceptibility of the alloy to Te-induced stress corrosion cracking in a molten salt environment. Understanding these factors is crucial for designing corrosion-resistant structural materials and optimizing the operating conditions of the MSR.