Corrosion behavior of ferritic FeCrAl alloys in simulated BWR normal water chemistry

Journal of Nuclear Materials 545 (2021) 152744

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Corrosion behavior of ferritic FeCrAl alloys in simulated BWR normal water chemistry Peng Wang a,∗, Slavica Grdanovska b, David M. Bartels b, Gary S. Was a a b

Department of Nuclear Engineering and Radiological Sciences, University of Michigan, Ann Arbor, MI, 48109, USA Notre Dame Radiation Laboratory, University of Notre Dame, Notre Dame, IN, 46556, USA

a r t i c l e

i n f o

Article history: Received 1 July 2020 Revised 5 December 2020 Accepted 14 December 2020 Available online 17 December 2020 Keywords: Steel Raman spectroscopy SEM STEM Higher temperature corrosion Ion irradiation

a b s t r a c t Ferritic iron-chromium-aluminum (FeCrAl) alloys with varying Cr content were subjected to in-situ irradiation-corrosion with energetic protons or electron beams in pure water with various oxygen concentrations. Oxide microstructure and composition were characterized to compare the corrosion under irradiation with the thermal corrosion in the same water conditions. The dominant effect seems to be the production of H2 O2 at the surface by the intense proton and electron beams. The corrosion potential of FeCrAl under BWR-NWC condition is somewhat above the Cr2 O3 /CrO4 2− phase boundary, oxidizing the Cr(III) to Cr(VI) in the outer part of the film, but with some Cr(III) remaining in the interior of the oxide. The beam increases the corrosion potential, significantly reducing Cr(III) content throughout the oxide. The crystallites on the surface, which form by re-precipitation from solution, do not contain Cr because at low potential Cr(III) is not soluble, and at high potential Cr(VI) is quite soluble.

1. Introduction Fukushima Daiichi nuclear accident in 2011 has redirected the world’s nuclear fuel R&D activities from waste management to enhancing the accident-tolerance of light water reactor (LWR) fuel [1]. Accident tolerant fuel (ATF) has been one of the focus areas over the past decade to provide a solution to the shortcomings of zirconium-based cladding materials under accident conditions. Zirconium alloys are used as fuel cladding materials in LWRs and undergo an exothermic reaction with steam at high temperatures encountered during a loss-of-coolant type accident that results in increased production of H2 due to its autocatalytic nature. The objective is to replace zirconium alloys with a material system with increased oxidation resistance in a high-temperature steam environment. Ferritic FeCrAl alloys can form an alumina film under accident conditions that will extend the coping time of fuel assembly in a loss of coolant accident (LOCA) event. However, these alloys must also perform equally well or better than the existing cladding material under normal operating conditions. Autoclave experiments confirm several candidate alloys that might perform very well, but the radiation field is omitted from these tests. A previous study of FeCrAl alloys by Wang et al. showed that proton and electron

∗

Corresponding author.

https://doi.org/10.1016/j.jnucmat.2020.152744 0022-3115/© 2020 Elsevier B.V. All rights reserved.

© 2020 Elsevier B.V. All rights reserved.

irradiation in a hydrogenated water environment (i.e., 3 wt. ppm dissolved hydrogen) increased the corrosion potential and resulted in oxidation of most of the Cr(III) in the film to Cr(VI) which is soluble [2]. Boiling water reactor conditions are more oxidizing than those in a PWR, so there is a need to evaluate the behavior of these alloys under these conditions as well. The performance of FeCrAl cladding material must be evaluated for accident tolerance fuel application in the BWR environment. However, the in-core exposures and the subsequent post-irradiation characterization would take a considerable amount of time and cannot satisfy the needs for a quick sample turnaround. In this paper, we report short-term accelerator-based in-situ irradiation-corrosion experiments intended to provide quick evaluation for candidate material screening and understand the simultaneous and individual effects of irradiation damage and water radiolysis on various FeCrAl alloys in BWR conditions. It should be noted immediately that the corroding alloys are probably not in a steady-state due to the limited length of these in-situ irradiation exposures. Therefore, the short-term study of in-situ irradiationcorrosion should be defined as an ’initial transient’ condition. The analysis was conducted on the oxide species, morphology, and alloying element re-distribution through oxide microstructure characterization.

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Journal of Nuclear Materials 545 (2021) 152744

at 2 μA/cm2 , resulting in a flux of 1.25 × 1013 protons/cm2 -s and a damage rate in the metal at the metal/water interface of 3 × 10−7 dpa/s. The linear energy transfer (LET) to the water was ~ 5.6 × 10−11 J/cm per proton, which results in a dose rate of ~ 400 kGy/s, assuming the initial energy of the proton entering the water was 1.4 MeV. The dose rate in this study is significantly higher than the maximum dose rate found in the LWR core region (~ 10 kGy/s) [7]. The sample temperature was measured during irradiation in a previous in-situ study of stainless steel 316L in 320°C water to verify that the beam heating effect was minimum [6]. Measurements were made using an infrared pyrometer on the vacuum-facing side of the sample. At a beam current density of 10 μA/cm2 , a 3°C temperature rise was measured. A much lower temperature rise for this analysis is expected at a proton current density of 2 μA/cm2 .

Table 1 Chemical compositions of the alloys used in this study. Alloy compositions are nominal compositions. Alloys

T35Y2 T54Y2 APMT

Compositions wt%

Irradiation Modes

Fe

Cr

Al

Mo

Y

Protons

Electrons

82.3 80.8 70.5

13.1 15.1 21.0

4.4 3.9 5.0

3.0

0.1 0.1 -

Y Y Y

Y Y Y

2. Experiment 2.1. Materials and sample preparation Three FeCrAl alloys were studied to evaluate their feasibility as ATF cladding materials in BWR water chemistry. All alloys have an equiaxed grain size of 30-50 μm. Model alloys T35Y2 and T54Y2 were manufactured via arc melting, casting, and rolling at Oak Ridge National Laboratory. Detailed microstructure characterization of these alloys can be found in Ref. [3]. Commercial Kanthal® APMT is a dispersion strengthened FeCrAl alloy fabricated through advanced powder metallurgy. The nominal chemical compositions of the alloys are shown in Table 1. Disc samples were fabricated out of materials using electric discharge machining (EDM). The as-machined discs have a 7.6 mm diameter and a starting thickness of 0.5 mm. The discs were mechanically ground and polished to a mirror finish using silicon carbide abrasive paper and diamond polishing solution. A final vibratory polishing procedure was introduced with a 50 nm colloidal-silica solution to remove any deformation layer. The final disc thickness was 80 ± 2 μm, determined based on the available beam energy and the thickness needed to ensure structural integrity throughout the experiment. The critical parameter for insitu proton-irradiation is the damage rate to the metal and the dose rate to the water at the metal/water interface. The damage profile through the metal is irrelevant in this case. Due to the much higher penetration depth of 2.5 MeV electrons in metals, the thickness of the electron irradiated samples was set at 250 ± 2 μm. A detailed analysis of proton penetration depth and sample thickness determination can be found in Ref. [2]. Disc samples were welded to the sample mount to form a watertight seal, using either pulse arc welding (proton irradiated samples) or electron beam welding (electron irradiated samples), respectively. A detailed description of the sample assembly can be found in Ref. [2,4].

2.3. Electron irradiation Electron irradiation was conducted on a 3 MV Van de Graaff accelerator at Notre Dame Radiation Laboratory (NDRL), using a 2.5 MeV electron beam. The current density was much higher than in the proton experiments, 120 μA/cm2 (i.e., electron flux at 7.5 × 1014 electron/cm2 -s), due to the much smaller LET of the electrons when entering the water (~3.04 × 10−11 J/cm per electron at ~ 1.5 MeV). The estimated dose rate from the electron beam is ~ 170 kGy/s. A higher water flow rate was required to remove the excess heat from the electron beam heating, to maintain the sample temperature at 288°C. The inlet and outlet of the autoclave temperature were monitored through two thermocouples inserted very close to the irradiated zone. Since the flow channel of the Hastelloy autoclave was narrow, the average between the inlet and outlet temperature was used as the temperature for the irradiated zone. A detailed description of the electron irradiation setup can be found in Ref. [2]. The autoclave at NDRL was made of nickelbase alloy Hastelloy C276. In addition to BWR-NWC water chemistry, Argon saturated water chemistry has been used with a dissolved oxygen level < 10 wt.ppb. 2.4. Post irradiation characterization Characterization of the surface oxide was conducted using transmission electron microscopy (TEM) on cross-section samples made by a focused ion beam (FIB) lift-out technique on a Thermo Fisher Helios 650 NanoLab SEM/FIB system. The straight crosssection was used for all FIB lift-out and TEM analyses. The oxide microstructure, thickness, and composition analysis were conducted on a JEOL 3100 R05. Bright-field (BF) and high-angle annular dark-field (HAADF) imaging techniques were used to obtain the oxide thickness data. Energy-dispersive X-ray spectroscopy (EDS) was used to obtain the composition information of the oxide layers and across the metal/oxide interface. Raman spectra were collected on a Renishaw inVia microscope with a RenCam CCD detector. The spectra were obtained using a 633 nm wavelength red laser at 50 mW laser power, with a dwell time of 900 s.

2.2. Proton irradiation In-situ proton irradiation-corrosion experiments were conducted at Michigan Ion Beam Laboratory (MIBL) using a pelletron particle accelerator with a terminal voltage of 2.7 MV, producing proton energies at 5.4 MeV. The sample was mounted in a miniature autoclave with a water circulation system, connected to a dedicated beamline of the accelerator. Water chemistry simulated the BWR-NWC condition with 2 wt.ppm of dissolved oxygen in pure water, with an electrical conductivity of 6.5 μS/m at room temperature. The entire water loop and autoclave were constructed out of SS316L. The FeCrAl disc sample served as the “window” to allow energetic projectiles (i.e., proton) to penetrate the sample thickness fully, and also as a pressure barrier between the high-temperature high-pressure water (288°C, 11 MPa) in the autoclave and the high vacuum (10−12 MPa or 10−7 torr) in the beamline. A detailed description of the experimental design and setup can be found elsewhere in Ref. [2,4–6]. The range of a 5.4 MeV proton in the water is approximately 100 μm after exiting the 80 μm thick sample. During the proton irradiation, the current density was kept

3. Results Oxide morphology, composition, and structure from both proton and electron irradiation-corrosion experiments are presented in the following sections. A complete, quantitative account of all samples tested is provided in the table at the end of the section. 3.1. General appearance Distinguishable regions were observed on the sample waterfacing surface of sample T35Y2 after in-situ testing under proton 2

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Fig. 1. The optical image, SEM image, and schematic of the various regions of sample T35Y2 after the in-situ proton irradiation-corrosion experiment at 288°C in NWC water (2 wt.ppm O2 ) for 24 hours. “IR” consists of a 1 mm diameter area that was directly proton irradiated, “UR” is the unirradiated region on the outer rim of the sample, and the “RAR” denotes the radiolysis affected region in between the UR and IR regions. Table 2 Three testing conditions for FeCrAl alloys in this study. Autoclave

Test conditions, 24 hours in-situ irradiation-corrosion

Tested Materials

UM 316L Autoclave ND Hastelloy Autoclave

Proton irradiation, high-purity water, 2 wt. ppm O2 , 288°C Electron irradiation, high-purity water, Argon saturated, 288°C Electron irradiation, high-purity water, 2 wt. ppm O2 , 288°C

All alloys in Table 1 T35Y2, APMT T54Y2

irradiation at 288°C in NWC water for 24 hours shown in Fig. 1. The short exposure time used in this study indeed only explored the initial stage of corrosion of these alloys. However, a mechanistic understanding of the radiolysis affected corrosion behavior can still be learned. The center irradiated region (IR) of the sample resembles the circular shape of the proton beam defined by an aperture behind the sample. Minimal beam heating (~5°C) of the sample and adjacent water was expected, which induced an upwards convective flow of water in front of the IR region. This convection flow also carries long-lived radiolysis products away from the IR region and forms a second area around the IR (the blueish color rim in the optical image in Fig. 1), termed the radiolysis affected region (RAR). The subtle color differences of the IR and RAR regions in the optical image is produced by the interference of light when reflected through different thicknesses of the oxide films. However, this small difference in oxide thickness was invisible under the SEM. The broader rim of the sample surface was categorized as the unirradiated region (UR), as it experienced neither the proton beam nor the concentrated water radiolysis products. The electron irradiated samples were similar in terms of surface appearances, though with an elongated RAR region extended further away from the IR region since the flow rate was much higher. 3.2. Oxide morphology and composition 3.2.1. Proton irradiation mode Table 2 lists the conditions used to study the material response under various oxygen concentrations. These should not be interpreted as corresponding to power plant conditions since plant conditions are far more complex than those used in this study. Fig. 2 shows the surface oxide morphology on each of the three samples, with images from each region following in-situ proton irradiationcorrosion experiments. All sample surfaces were fully covered with oxide particles, with particle size trending down from UR to IR region. The outer crystallite particles are formed by re-precipitation from the solution and contribute little to passivity. However, the morphology of the outer crystallites provides evidence of their local environment and formation process: undisturbed faceted crystals in the unirradiated region vs. rounded non-faceted crystals in the irradiated and radiolysis affected regions. There is also a tendency for increasing oxide particle size with increasing metal Crcontent, ranging from ~100 nm on low-Cr alloys to a few μm on

Fig. 2. SEM images of the oxide film formed on FeCrAl alloys after in-situ irradiation-corrosion exposure using a 5.4 MeV proton beam in NWC water (2 wt.ppm O2 ) at 288˚C for 24 hours.

high-Cr samples. Most of the oxide particles were packed randomly with a rounded crystal appearance. Raman spectra were collected on all samples at various regions to understand the effect of radiation on oxide species formation. Since the Raman laser has a limited penetration depth, the spectra typically reflect the response of the surface. Due to the extensive coverage of the surface oxide, the Raman responses were likely to represent the top layer oxide. Fig. 3 shows an example of the Raman responses of the oxide collected on T35Y2 after proton irradiation. All the peaks obtained in Raman wavenumber ranging from 200-750 cm−1 represent the hematite (α -Fe2 O3 ) responses, which are in good agreement with the literature data found in Ref. [8]. 3

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Journal of Nuclear Materials 545 (2021) 152744

Table 3 Collective data for FeCrAl alloy. Proton and electron irradiation data are marked (p) and (e), respectively.

Unirradiated Region (UR)

Outer Oxide

Oxide phasea Cr, Ni contentsb

Irradiated Region (IR)

Radiolysis Affected Region (RAR)

Outer Oxide

Oxide phase

Outer Oxide

Cr, Ni contents Oxide phase Cr, Ni contents

a b

Alloy/(Cr wt%) T35Y2 (13.1)

T54Y2 (15.1)

APMT (21)

α -Fe2 O3 (p) α -Fe2 O3 (e)

α -Fe2 O3 (p) α -Fe2 O3 & NiFe2 O4 (e)

α -Fe2 O3 (p) α -Fe2 O3 (e)

~6 at.%Cr(p) ~10 at.%Ni(e) α -Fe2 O3 (p) NiO, NiFe2 O4 (e) - (p) 5-15 at.%Ni(e) α -Fe2 O3 (p) α -Fe2 O3 (e) ~ 2 at.%Cr (p) ~12 at.%Ni (e)

~9 at.%Cr(p) ~10 at.%Ni(e) α -Fe2 O3 (p) NiO & NiFe2 O4 (e) ~ 2 at %Cr(p) ~5-40 at.%Ni(e) α -Fe2 O3 (p) α -Fe2 O3 & NiFe2 O4 (e) ~ 5 at.%Cr(p) ~5-40 at.%Ni(e)

~12 at.%Cr(p) ~10 at.%Ni(e) α -Fe2 O3 (p) NiO, NiFe2 O4 (e) ~ 2 at.%Cr(p) ~7 at.%Ni(e) α -Fe2 O3 (p) α -Fe2 O3 (e) ~ 10 at.%Cr(p) ~5-10 at.%Ni (e)

Oxide phases were determined using Raman spectroscopy, and SAED in TEM mode. Compositional analyses were determined using EDS in STEM mode.

Fig. 3. Raman spectroscopy data collected on T35Y2 samples after in-situ irradiation-corrosion exposures using a 5.4 MeV proton beam in NWC water (2 wt. ppm O2 ) at 288°C for 24 hours.

Fig. 4. STEM HAADF images of the metal/oxide interface of alloys after in-situ irradiation-corrosion exposures using a 5.4 MeV proton beam in NWC water (2 wt.ppm O2 ) at 288°C for 24 hours.

The Raman responses from all three regions of the irradiated samples identified α -Fe2 O3 as the surface oxide phase. This observation also holds for all other alloys tested under proton irradiation, as the oxide phase results indicated in Table 3. STEM/HAADF imaging was used to characterize the metal/oxide interface of the samples post-test. Fig. 4 shows the collective images of the metal/oxide interfaces of all samples. The images are orientated such that the dark-colored oxides were positioned in the center, separating the post-test Pt deposition layer (upper portion) applied to protect the underlying metal during FIB milling. The thickness of the oxides formed on these samples ranged from tens of nm to ~200 nm. In general, the oxide morphology in the UR regions agrees with the observation from the SEM characterization, where large oxide particles were seen to penetrate the metal and created a wavy metal/oxide interface. In Fig. 4a to 4c, the oxide layers on all alloys consist of a single layer oxide made of large single crystals. This is also evidenced by the unchanged oxide contrast throughout its thickness in these STEM images taken in HAADF mode, providing better Z-contrast than BF imaging mode. Fig. 4a and 4b show that the oxide particle sizes in the RAR regions of T35Y2 and T54Y2 are smaller than in the UR region, while the opposite is the case for APMT, Fig. 4c. This behavior is consis-

tent with the SEM observation in Fig. 2. The oxides found in the RAR region are all single-layer oxides. The oxides are composed of individual crystallites in Fig. 4a to c, re-precipitated from the solution. The IR region oxides are consistently thinner than those in the UR and RAR regions on all samples, and the IR region oxides show consistently smaller crystallite sizes and less undulation at the metal/oxide interface. The overall trend for oxide thickness evolution on these alloys is evident. As the concentration of water radiolysis products increases from the UR region to the IR region, the resulting oxide thickness decreases. This is unlikely to be a temperature difference between the UR and IR regions since the beam heating effect was less than 3°C, which is insignificant to cause a shift in the corrosion rate. As a result, the large water chemistry change was thought to be responsible for the oxide thickness differences. From an alloy Crcontent perspective, as the alloy Cr level increases, the oxide thickness decreases. In Fig. 5, EDS line-scan profiles are presented as a function of alloy composition and region. In the UR region, the Cr concentra4

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Journal of Nuclear Materials 545 (2021) 152744

Fig. 5. EDS line-scan profiles across the metal/oxide interface of alloys after in-situ irradiation-corrosion exposures using a 5.4 MeV proton beam in NWC water (2 wt.ppm O2 ) at 288°C for 24 hours. Oxide thicknesses shown in the EDS plot do not represent the average oxide thickness. For accurate oxide thickness measurement, please refer to Fig. 10.

tions in the oxide are similar for T54Y2 and APMT, ~ 10 at.% Cr. For T35Y2, the Cr content in the oxide varies from 5-15 at%. The metal (Fe+Cr) to oxygen ratio suggests the oxide stoichiometry is M2 O3 , except for T35Y2, where the metal to oxygen ratio is close to 1. However, the Raman response suggests the oxide is α -Fe2 O3 in nature. This means that all the outer crystallites formed in the UR region have the hematite characteristics with some Cr content. In the RAR region, although Cr is still being detected in the oxide, its concentration has dropped from 9±3 at. % in the UR region to 6±4 at. % in the RAR region. The oxide stoichiometry remains at O/M ~ 3/2, consistent with hematite. In the IR region, the Cr content in the oxides is further reduced to < 2 at. %, which makes the oxide composition close to hematite. As the concentration of water radiolysis product increases across UR, RAR, and IR regions, there is a general trend of decreasing Cr-content in the oxide, and the oxide composition shifts from a Crx Fe2-x O3 towards α -Fe2 O3 . At low potential, e.g., UR region, Cr(III) is not soluble, it can only be moved through the oxide crystallites via solid-state diffusion. At high potentials, e.g., RAR and IR region, CrO4 2− is soluble, and it does not participate in the re-precipitation process. As the alloy Cr-content increases, more Cr can be found in the resulting oxide. There is also a noticeable reduction in oxide thickness in IR regions compared to the UR region across the alloys. More importantly, the protective Cr-rich oxide typically observed near the metal-oxide interface following long-term exposures was not apparent in these 24-hour exposures.

3.2.2. Electron irradiation mode The alloys were also tested under similar water conditions using electron irradiation. Different water chemistries were also employed to determine the alloy response/sensitivity to different dissolved oxygen levels. Two water chemistries are listed in Table 2, Argon saturated water and NWC. Fig. 6 shows the morphology of the surface oxides formed on electron irradiated samples, categorized as a function of irradiation regions and alloy composition. Under UR conditions, the oxide particles formed in Argon saturated water (i.e., T35Y2 and APMT) show a mixture of faceted particles of all sizes. The particle size on T35Y2 is larger and more faceted than it was on APMT. The oxide particles found on T54Y2, which was exposed in NWC, can be categorized into two groups, with one group having bulky and isolated particles and the other group consisting of small and densely packed particles. The oxide particle in RAR regions found on T35Y2 and APMT look identical to the UR cases. There are fewer large particles on T54Y2 in the RAR region; the overall particle density stays the same. In the IR region, there is a significant reduction in oxide particle sizes on T35Y2 and APMT, as well as T54Y2. In all cases, there is a trend of reduction of particle size across UR, RAR, and IR regions as water radiolysis product increases. Raman spectra were collected on all three samples at various regions to determine the oxide phases concerning their local water chemistries. Fig. 7 shows an example of the Raman responses collected from the oxides on T35Y2 after electron irradiation in Argon saturated water. The Raman spectra from UR and RAR regions are 5

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Journal of Nuclear Materials 545 (2021) 152744

Fig. 6. SEM images of the oxide film formed on FeCrAl alloys after in-situ irradiation-corrosion exposures using a 2.5 MeV electron beam in NWC water (2 wt.ppm O2 ) and Argon saturated water at 288°C for 24 hours.

environment, including Ni2+ from corrosion of the Hastelloy autoclave. Fig. 8 shows the collective STEM images of the metal/oxide interfaces of all samples irradiated with electrons. In the Argon saturated water exposed cases, T35Y2 and APMT, the oxide thickness is typically in the range of ~200 nm across all three regions. The oxide layers consist of large oxide crystals on the outer part of the layer and smaller crystals close to the wavy metal/oxide interface. There are only two minor differences seen on T35Y2 and APMT. First, the size of the outer oxide particles decreases as the alloy Crcontent increases. Second, there are also noticeable gaps and voids between the larger oxide crystals, and the local oxide penetration seems to be deeper below these features. Fig. 9 shows EDS line-scan profiles collected from electron irradiated alloys as a function of alloy composition and water conditions. In cases where T35Y2 and APMT were exposed to argon saturated water, the UR regions showed consistent oxide results with Nix Fe2-x O3 composition, where x is about 0.5. The oxide Ni concentration increases to x = 0.66 in the RAR regions, with oxygen level at a steady ~ 60 at. %. In the IR region, the Ni content in the oxide is further reduced to ~ 6±1 at. % and gradually decreases towards the metal/oxide interface. In NWC exposed T54Y2, the oxide composition in the UR region is similar to a NiFe2 O4 composition with minor metal deficiency. As the concentration of water radiolysis products increases, the outer part of the oxide has a higher amount of Ni incorporated into the oxide, accompanied by a lower oxygen concentration in the RAR region. This Ni-enriched layer becomes more pronounced in the IR region than the UR region, where the Ni concentration increases to 40 at. % Ni. In contrast, NWC water seems to impede the general formation of large outer oxide crystals on T54Y2; only large crystals were observed in the SEM image in Fig. 6, especially in UR and RAR regions. However, the TEM lamellas prepared for T54Y2 did not capture these large crystals. The oxide consists of finer crystals on T54Y2 in the UR and RAR region, with an average thickness of ~ 100 nm. The oxide layer in the IR region is slightly thicker and consists of larger oxide crystals. The electron irradiation shows a negligible impact on oxide morphology/oxide thickness for alloys (i.e., T35Y2 and APMT) exposed in the Argon saturated water, while noticeable oxide thickening was observed on T54Y2 for NWC exposure. The Argon saturated water exposure produced a thicker outer oxide, and particle sizes were much larger than exposures conducted in NWC water. All of the FeCrAl alloys tested in the electron irradiated cases have some Ni incorporated in their oxides. In Argon saturated wa-

Fig. 7. Raman spectrum was collected on T35Y2 after in-situ irradiation-corrosion exposures using a 2.5 MeV electron beam in Argon saturated water at 288°C for 24 hours.

typical of an α -Fe2 O3 oxide, while the IR region shows the spectrum from a NiFe2 O4 spinel oxide. Since Ni was not the alloying element of FeCrAl, the Ni content found in the oxide was unintentionally introduced from the Hastelloy 276 of the autoclave, a highnickel alloy which is generally recommended for high-temperature water applications in oxidizing condition. It should be noted that nickel from the SS-316L autoclave was not observed on samples in the proton irradiation experiments. The difference may be primarily due to the much higher flow velocity of the water in the integral preheater loop of the Hastelloy autoclave. APMT sample exposed in the same water chemistry developed a Raman spectrum very similar to T35Y2. On T54Y2 exposed to NWC, the UR and RAR spectra showed characteristics of both α -Fe2 O3 and NiFe2 O4 spinel oxides, whereas only NiFe2 O4 spinel was observed in the IR region. Note that a 568 cm−1 peak is present in the IR regions of all alloys. This response is distinct from the NiFe2 O4 response, and it is identified as possibly NiO [9]. The water chemistry differences influence the oxide species in the UR and RAR regions, whereas the IR regions consistently show NiO and NiFe2 O4 signals. Given that none of the alloys under study contain nickel, it is clear that the surface oxides result from the precipitation of metal ions from the 6

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Journal of Nuclear Materials 545 (2021) 152744

Fig. 8. STEM HAADF images of the metal/oxide interface of alloys after in-situ irradiation-corrosion exposures using a 2.5 MeV electron beam in NWC water (2 wt.ppm O2 ) and Argon saturated water at 288°C for 24 hours.

Fig. 9. EDS line-scan profiles across the metal/oxide interface of alloys exposed in argon saturated water (i.e., T35Y2 and APMT) and NWC water (i.e., T54Y2) condition under 2.5 MeV electron irradiation at 288°C for 24 hours. Oxide thicknesses shown in the EDS plot do not represent the average oxide thickness. For accurate oxide thickness measurement, please refer to Fig. 10.

ter, the Ni distribution within the oxide is very uniform. In NWC water, there is a gradient of Ni content in the developing oxide of the RAR and IR regions, and a Ni-enriched layer was observed near the surface. The enriched Cr oxide typically observed near the metal-oxide interface after long exposure [10,11] was not observed in these 24-hour exposures in both water environments.

Based on previous studies, mass losses are expected to be much smaller under the NWC condition than corrosion under hydrogenated water chemistry [2]. Indeed, a long-term autoclave exposure study conducted on FeCrAl alloys by ORNL in NWC water demonstrated that the corrosion rate differences are insignificant in a comparison between the alloys [10]. As illustrated in Fig. 10, the UR regions on both proton and electron irradiated T35Y2 and T54Y2 yield an oxide thickness close to the C35M3 (an ORNL developed FeCrAl alloy with 13 wt. % Cr) reference case. It should be noted that C35M3 is compositionally similar to T35Y2 and was exposed in the NWC (1 wt.ppm dissolved oxygen) for over nine months, and the oxide thickness remained constant after 3-month exposure [10]. Unlike the corrosion behavior in the hydrogenated water chemistry, in which the oxide thickness increased over time, C35M3 continues to show a constant weight and consistent oxide thickness over nine months when exposed to NWC. A Cr-enriched inner layer, was clearly visible on C35M3 at the metal/oxide interface after exposure. However, in our study, the Cr-enriched inner oxide was not observed on either T35Y2 or T54Y2. Presumably, the exposure time was too short to accumulate Cr at the metal/oxide interface, even though these UR samples show oxide thickness similar to the C35M3 after just 24 hours. At steady-state, the dynamic Fe dissolution and re-precipitation of outer oxide from solution are in balance, resulting in a constant oxide thickness with time. In contrast to these alloys, APMT over long-term exposure showed minor weight gain and thicker oxide that can be inter-

4. Discussion 4.1. Corrosion behavior without irradiation There have been many autoclave studies of stainless steel corrosion in BWR “normal water chemistry” relevant to nuclear reactor cooling loops [10–17]. Typically results are reported for the passive film characteristics established after several hundred hours of exposure or longer. Raiman et al. [10] have reported several representative FeCrAl alloys following nine-month exposure to water with 1 wt.ppm oxygen. From long-term autoclave exposure studies on FeCrAl alloys [10,11], it has been reported that the overall oxide structure was a double-layer oxide in both hydrogenated or oxygenated water. The double-layer oxide consisted of a Fe-rich outer oxide and a thin Cr-rich inner oxide at the metal/oxide interface. A lower mass loss was also observed for these FeCrAl alloys under oxygenated conditions (1,0 0 0 wt.ppb O2 ) compared to the hydrogenated condition (63 wt.ppb H2 ) at 290°C [10]. Therefore, it is interesting to compare these results with the characteristics of our UR samples, which result from only 24 hours of exposure. 7

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Journal of Nuclear Materials 545 (2021) 152744

Fig. 10. Box plot showing oxide thickness of FeCrAl alloys exposed for 24 hours in proton and electron irradiated experiments compared to literature results in a long-term (up to 9-month) autoclave exposure [10].

preted as a balanced higher Fe dissolution due to faster corrosion of the alloy and weight gain from the re-precipitated outer oxides [10]. Compared to C35M3, a thinner and more Cr-enriched inner oxide was observed at the metal/oxide interface of the long-term exposed APMT, accompanied by a significantly thicker (~ 400 nm) outer oxide, also noticeable in Fig. 10. However, there are no significant Cr-enrichment in the UR regions of our proton and electron irradiated APMT samples, again probably due to the short exposure time. The UR oxide thicknesses of the APMT in this study (shown in Fig. 10) after 24 hours of exposure are still far from the steadystate conditions. ECP is measured as the voltage difference between a metal immersed in a given environment (i.e., an electrolyte) and a standard reference electrode (e.g., standard hydrogen electrode, SHE). It reflects the average potential energy (or Fermi level) of the electrons in the material at the surface, which results from the balance of the total cathodic and anodic reaction currents. It is a response of the corroding surface to the electron transfer reaction rates, which are highly influenced by the reactant concentrations, surface composition, the thickness of the oxide and even the flow velocity of the liquid. When comparing the corrosion behavior of materials, it is essential to compare systems that are dominated by the same chemical reaction after passivation; in this case, the corrosion of steel is dominated by the Fe→ Fe2+ reaction. Hence, it is appropriate to compare ECP of FeCrAl alloys with the other stainless steels (e.g., SS304 and SS316L) that were used in the same reactor environment. Open circuit ECP of stainless steel samples (e.g., type 304 SS) in the NWC condition has been reported by numerous researchers, ranging from -100 to 100 mVSHE for a dissolved oxygen level of 200 wt.ppb to 8 wt.ppm [16,18]. After an initial induction period to establish the (hematite) oxide surface, the corrosion potential for 2 wt.ppm dissolved oxygen remains relatively constant, near 0.0 mVSHE [18]. In the absence of radiation, deaerated water (ie., Argon saturated water, < 10 wt. ppb O2 ) would give an ECP on 288°C stainless steel of < -350 mVSHE [16]. An E-pH diagram for the Fe-Cr-O system under NWC conditions is shown in Fig. 11a. The lines separating oxidation states stand for conditions of equal thermodynamic stabilities of the two states in equilibrium (assuming 1 × 10−6 molar activity of the solution phase species). It should be understood that the phases are always in equilibrium, and both forms may be observed under conditions “near” the boundary lines. Due to the relatively positive ECP voltage imposed by the NWC experimental conditions, hematite (α -Fe2 O3 ) will be overwhelmingly preferred over the spinels, e.g., magnetite (Fe3 O4 ) and chromite (FeCr2 O4 ).

Hematite outer precipitates were the main oxide species detected on the FeCrAl alloys exposed in the ORNL long-term corrosion study. Up to 12 at.% Cr (or 18 wt.%) was detected in the APMT outer re-precipitated hematite crystallites, whereas the precipitates found on C35M3 only contains ~ 2 at.% Cr [10]. Consistently, the same amount of Cr has also been detected in the UR region precipitates on APMT, whereas the T35Y2 only has 6 at.% in its outer oxide, as shown in Fig. 5. Outer hematite oxide precipitates containing some Cr has been reported on autoclave exposed SS316/304 under NWC condition [19–21]. The enrichment of Cr in the outer oxide cannot be due to the precipitation of Cr from solution, because the NWC condition would favor the highly soluble CrO4 2- over Cr(III). The preferential diffusion of Cr from the underlying matrix to the outer part of the oxide is also unlikely, since Cr has been proved to have a much slower diffusion rate than Fe in the oxides, and Cr depletion was not detected in the metal near the metal/oxide interface in Fig. 5. It is proposed by Kuang et al., that at the initial stage of corrosion, selective Fe dissolution would result in a Cr enriched layer in the metal, followed by solid-state reactions with oxygen to form the first layer of Cr-enriched hematite nuclei [21]. Subsequent hematite growth was through precipitation of metal ions (e.g., mainly Fe) from solution. The alloy Cr content also closely influences the build-up of the Cr at the metal/oxide interface. However, due to the short exposure time of this study, the accumulation of Cr at the metal/oxide interface had not occurred yet. The Crx Fe2-x O3 oxide crystals are bulky; they only offer partial protection of the metal substrate. The uncovered regions between two adjacent oxide particles are likely to be attacked, evidenced by the large undulation of the metal/oxide interface in Fig. 4a and e. The degraded protectiveness of this Crx Fe2-x O3 oxide has been reported by Asteman et al. compared to its single constituents, α Fe2 O3 , or Cr2 O3 [30]. The C35M3 sample of Raiman, et al., showed constant oxide layer thickness over 9-month of exposure while continuously losing weight in NWC water [10]. This observation also suggests the Crx Fe2-x O3 oxide is less corrosion resistant. Since the Hastelloy autoclave used in the electron irradiation experiments was releasing nickel, oxide phases observed in these experiments followed the prediction in Fig. 11b for the Fe-Ni-CrO system. Without direct irradiation, a hematite-type Nix Fe2-x O3 oxide was precipitated under Ar-saturated water conditions, while the NiFe2 O4 phase is the dominating oxide in the NWC condition due to the more positive ECP. In the Argon saturated case, the corrosion potential of FeCrAl is controlled by the H2 -H2 O line, which gives a lower corrosion potential than the NWC condition. The hematite-type oxide observed in Fig. 7 also reflects the lower cor8

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Journal of Nuclear Materials 545 (2021) 152744

Fig. 11. E-pH diagrams of (a) Fe-Cr-O system and (b) Fe-Cr-Ni-O system, representing the BWR-NWC environment at 288˚C. The molality of the solute is set at 1μM/kg. The red dot indicates electrochemical potential at the starting condition (UR). (For interpretation of the reference to color in this figure legend, the viewer is referred to the web version of this article.) The plot is after references [12,22–29].

rosion potential observed in the UR and RAR regions of the T35Y2 and APMT samples. Beverskog and Puigdomenech have also observed the α -Fe2 O3 formation in the aqueous system of Fe-Cr-Ni [31] when excess Fe is present in the water with soluble Fe/Ni ratio > 2. This is due to the suppression of Ni release under deaerated water condition (i.e., Argon saturated) with an estimated ECP of -350 mVSHE ) [16], while enhanced Fe corrosion is expected. The corundum hematite-type Nix Fe2-x O3 oxide was similar to the Crx Fe2-x O3 oxide found on proton irradiated samples, which produced the same Raman response despite the different chemical compositions.

boundary (i.e., ~ -0.47 VSHE ) [17]. Because this ECP is more positive than the oxidation potential of Cr(III) to Cr(VI) transition at 300°C ~ 0 mVSHE [28], it is expected there will be some Cr dissolution under these conditions. In the proton irradiation experiment reported here, an LET of ~ 390 MeV/cm in the water (density of 0.74 g/cm3 ) can be expected from a 1.4 MeV proton (remaining energy when exiting from FeCrAl foil). All of the energy is deposited into the water within microns of the surface, leading us to estimate a near-surface water dose rate of ~ 400 kGy/s for the current density of the experiment. Electrons with similar energy would only have an LET rate of 1.9 MeV/cm, resulting in a dose rate of 170 kGy/s at the much higher current density used. Closer matching of the dose rates to the proton irradiation was not possible because of significant sample outgassing in the e-beam experiment. Both electrons and positive ions interact with the water mainly by Coulombic perturbation, exciting all of the dipole-allowed transitions of the liquid. By far, the greatest portion of the energy produces ionization and electronic excitation. The result of this interaction is the production of radiolysis products that include ionic and radical species along the path of travel: · e− aq , · HO, · H, · O2 − , · HO2 , H3 O+ , OH− , H2 O2 , and H2 [36]. Because of the greater track densities (corresponding to the LET difference), proton radiolysis will favor more H2 O2 and H2 production from prompt recombination of the radicals. Even with a perfect matching of the average dose rate, the resulting steadystate concentration of water radiolysis products (reached within milliseconds under both irradiation modes) would be different in the two experiments due to the differences in primary radiolysis production rate (G-values) of species created by protons and electrons at the same energy. The concentrations of radiolysis products produced in these experiments were simulated using the G-values and reaction rate constants at 288°C, published in the review of Elliot and Bartels [37]. In the simulation, radiation is “turned on” for 0.2 seconds and then “turned off”. Steady-state is reached within milliseconds, allowing qualitative identification of the “beam on” condition with the IR region of the samples and the subsequent “beam off” condition with the RAR region. The comparison is only meant to be qualitative as the model only represents homogeneous water and includes no interaction with the oxide surface or diffusion out of

4.2. Effect of radiolysis on corrosion potential Although the autoclave exposures answered some of the uncertainties regarding FeCrAl corrosion behavior in a non-irradiation environment, the normal operating condition of a BWR is more complicated than an autoclave test can simulate. During in-core residence, the cladding will experience radiation damage to the metal lattice and the evolving oxide. The corrosion behavior under the influence of water radiolysis products may differ from that in dissolved oxygen, even though the resulting ECPs are similar [32]. In boiling water reactors (BWRs), water at high temperature (about 560 K) and of high purity (electrical conductivity less than 11 μS/m [33] at room temperature, with the presence of radiolytically produced oxygen, hydrogen peroxide, and hydrogen) is circulated through the reactor core. Oxygen and hydrogen peroxide are produced by water radiolysis in BWRs under normal water chemistry (NWC) conditions [32,34]. The corrosion behavior of stainless steels in high-temperature water containing hydrogen peroxide was different from that in the coolant with only dissolved oxygen [16,32]. For the same dissolved concentration, the ECP of the stainless steel is more positive for the H2 O2 containing water compared to the water that contains O2 [16]. Hydrogen is often intentionally added to feedwater in some plants to mitigate stress corrosion cracking of structural materials. Without the addition of dissolved hydrogen, the oxygen level of the BWR coolant is typically at 100 wt.ppb, and hydrogen peroxide is at 200 wt.ppb in the recirculation loop [35]. The resulting ECP could be in the range of 0-100 mVSHE , which is much higher than the ECP for the magnetite (Fe3 O4 ) to hematite (α -Fe2 O3 ) phase 9

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Fig. 12. Simulation of water radiolysis for 0.2 seconds of irradiation with 1.4 MeV protons and relativistic electrons at 288˚C, in NWC and argon-saturated water chemistries. Steady-state concentrations are reached within milliseconds. The simulation result presented here should be considered as semi-quantitative only, because proton radiolysis product yields have not been measured experimentally at high temperatures. Table 4 Electrochemical reaction parameters at elevated temperature. Electrochemical reaction Cathodic H2 O2 + e– → OH + OH– (Eq.1) O2 + 2H2 O + 2e– → 2O2 – + 2H2 (Eq.2) Anodic H2 O2 →O2 + 2H+ + 2e– (Eq.3) Fe → Fe2+ + 2e– (Eq.4)

Standard electrode potential (VSHE )

Electrochemical reaction rate constant (m·s−1 )

-0.04 -0.04

5.0 × 10−5 [42] 2.0 × 10−7 [42]

-0.48 -0.44

10−3 [34,42] 10−2 [42]

the radiation zone. Fig. 12 shows the results of the three conditions used in the study. Even in Ar-saturated water, the steady-state concentrations of H2 , O2 , and H2 O2 were calculated to be orders of magnitude greater than those of any other radiolysis species, and therefore only these need to be considered. The concentration of O2 in NWC water is ~ 20 0 0 wt.ppb (62 micromolar), and this remains relatively constant [38,39]. The most significant effect of radiolysis is the production of hydrogen peroxide (H2 O2 ). The longlived H2 O2 can have a large impact on water chemistry locally in the IR region, and can be carried away by convective flow to impact the corrosion potential in the RAR region [40]. The metal oxide/water interface in this study is highly heterogeneous and known to be catalytic for H2 O2 decomposition due to the presence of α -Fe2 O3 on all metal surfaces [41]. The precipitated α -Fe2 O3 crystal surfaces have large surface areas and can act as catalytic sites to decompose H2 O2 into water and O2 . Typical cathodic and anodic reactions occurring on passivated stainless steel in NWC water are listed in Table 4:. Uchida and coworkers have demonstrated that hydrogen peroxide under NWC conditions will tend to dominate both the anodic and cathodic partial currents on passivated stainless steel for H2 O2 concentrations over 5 μM [12,18,26]. In the presence of this much H2 O2 , O2 addition up to 200 wt.ppb (6 μM) makes virtually no change to the ECP [18], which implies the large ratio of reduction rates for H2 O2 vs. O2 shown in Table 4 [42]1 . As H2 O2 concentration increases, the ratio of anodic and cathodic partial currents from H2 O2 oxidation or reduction remains constant, almost independent of [H2 O2 ] over a wide range of [H2 O2 ]. As a result, the corresponding ECP also remains nearly constant independent of [H2 O2 ]. On a fresh SS304 sample with H2 O2 present, the ECP

is near -0.2VSHE . As the oxide grows thicker, the (anodic) release of metal ions becomes slower, and after 30 hours, the ECP settles near 0.0 VSHE, where it is completely controlled by H2 O2 reactions [12,42]. Radiolysis produces roughly equal amounts of H2 and H2 O2 , and an important question is how much the oxidation of hydrogen on stainless steel contributes to the partial anodic current. According to the work of Uchida et al., H2 O2 still dominates the stainless steel corrosion potential under BWR-HWC conditions ([H2 ] < 50 μM) [12], which implies that the H2 O2 oxidation rate constant is also much higher than the hydrogen oxidation rate constant on this surface. The [H2 ] added to the water only moves the H2 -H2 O line vertically downwards on a Pourbaix diagram. However, this decrease of corrosion potential is small in magnitude compared to the large increase of potential by adding O2 /H2 O2, as indicated by Andresen in Ref. [38]. The radiolysis calculation of Fig. 12 suggests a steady-state concentration of 300 μM (~ 10 wt.ppm at 400 kGy/s) of H2 O2 was produced under the NWC condition by the proton irradiation. Given the large excess of H2 O2 over O2 in the proton experiment under the NWC condition, the H2 O2 will dominate the partial cathodic current. Concurrently, a nearly equal amount of H2 is also produced in the proton experiment. However, the dominating effect of H2 O2 on the anodic current observed in the case of HWC (previous paragraph) should also be expected in this case, as indicated by the model of Uchida, et al [12]. The steady-state ECP should settle near 0.0 VSHE or higher. The oxides we observe in the IR region are dominated by hematite, while the UR and RAR oxides are dominated by hematite-type Crx Fe2-x O3 . Based on the Pourbaix diagram of Fig. 11a, these observations would be consistent with a shift of the corrosion potential to a more positive value in the presence of the large H2 O2 concentration, favoring the formation of soluble CrO4 2− . The thickness of the protective oxide layer also affects the corrosion potential. It is unclear whether the IR oxide thickness represents the final passive layer for this [H2 O2 ] condition or

1 The numbers of Table 4 have been taken from Table 2 of reference [42], a conference proceedings paper that is nearly identical to the subsequent publication in Ref. [12]. However, the Table 2 of Ref. [12] contains different reaction rate numbers that appear to be generic guesses from previous decades. We have confirmed with Dr. Uchida that the Ref [12] table is in error. Nearly equal reduction rates for H2 O2 and O2 cannot explain the ECP observations of Ref. [18] and are inconsistent with the explanations given in Ref. [12].

10

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Journal of Nuclear Materials 545 (2021) 152744

if the system has simply not reached the final passive state after 24 hours. In the e-beam experiment in NWC, [H2 O2 ] is predicted to be at 15 μM, and [H2 ] near 22 μM at a dose rate of 170 kGy/s. Here again, the behavior of the corrosion potential should be very similar to the proton irradiation case described above, and the H2 O2 will be dominating the cathodic and anodic reactions. However, due to the large Ni concentration from the corrosion of the Hastelloy autoclave, the stable oxide phase under the NWC condition is NiFe2 O4 . In Fe-Cr-Ni systems, as shown in Fig. 11b, if there was excess Ni in the water, i.e., Fe/Ni<2, a combination of NiFe2 O4 and NiO was expected [28,31]. Indeed, NiFe2 O4 and NiO were detected in all regions of T54Y2 and all the IR regions of the electron irradiated samples, indicating that soluble Fe ions were overwhelmed by the Ni ions in the water due to corrosion release from the Hastelloy autoclave. Finally, in the e-beam Ar-saturation experiments, [H2 O2 ] is expected to be near 10 μM, [H2 ] near 25 μM, and [O2 ] near 6 μM at a dose rate of 170 kGy/s. The calculated [H2 O2 ], [H2 ], and [O2 ] concentrations appear to fit the criteria of Uchida and coworkers for the BWR hydrogen water chemistry, where H2 O2 still controlled both the cathodic and anodic partial currents [12], and a corrosion potential near 0.0 mVSHE is to be expected. The experiment confirms this expectation as the oxide phase changes from α -Fe2 O3 in the UR oxide to the NiFe2 O4 phase in the IR oxide on the electron irradiated samples (i.e., T35Y2 and APMT) in Ar saturated water.

of the alloy, resulting in a higher Fe release rate. Since all reactors were constructed out of materials with Ni as a common alloying element, large dissolved Ni concentrations in the BWRs coolant are unavoidable during reactor operation. It has been demonstrated in the electron irradiation experiments that high Ni content in the water favors the formation of nickel ferrite on the FeCrAl alloys, and NiFe2 O4 was also proven to suppress cobalt deposition and overall corrosion rate of stainless steel components [44]. However, if excess Fe is present in the water, hematite may be the dominating oxide species, which offers less protection over nickel ferrite. As a result, controlling the Fe concentration in the coolant becomes vital to optimize the corrosion resistance of these FeCrAl alloys in an oxygenated environment. 5. Conclusions This study aimed to evaluate the potential effect of a radiation environment on the corrosion response of FeCrAl alloys. The 24hour exposure time is likely insufficient to reach a steady-state but may be described as an “initial transient” state. Alloys with Cr content from 13 to 21% were tested. The following conclusions can be drawn from the results and analysis in this paper: •

4.3. Applicability to BWR conditions

•

This study only addresses the ‘initial transient’ corrosion response of these FeCrAl alloys since the exposure time is very short. However, during reactor startup after refueling, the dissolved oxygen and hydrogen peroxide concentrations in the reactor coolant would be relatively high, typically at 10 0 0 – 20 0 0 wt.ppb at the core outlet region, producing a strongly oxidizing condition similar to the conditions employed in this study. Therefore, the results obtained in this study could be representative of some specific conditions during the BWR operation, especially when a new fuel assembly of FeCrAl cladding is first introduced into the reactor coolant. Although ferrous ion oxidation and Cr dissolution are the main processes in the corrosion of FeCrAl in an oxygenated environment, the overall corrosion rate of the alloys is expected to be relatively low. This has been demonstrated by some long-term autoclave exposures conducted on FeCrAl variants at ORNL and GE, where FeCrAl displayed a very low corrosion rate in NWC, almost an order of magnitude lower than HWC [10,11]. Under oxygenated water, the FeCrAl alloys behaved similarly to stainless steels and nickel alloys containing chromium. However, for the past 60 years, Cr-containing stainless steels and nickel alloys performed well in the reactor coolant environment at nearly 300°C with a smaller concentration of oxygen (e.g., < 200 wt. ppb) controlled by hydrogen injection into the coolant, resulting in an ECP < -230 mVSHE . The critical element for their corrosion resistance can be attributed to an inner FeCr2 O4 or Cr2 O3 oxide layer and an outer NiFe2 O4 or NiO oxide [43]. This seems to be aligned well with the results from the electron irradiated FeCrAl. The main concern under irradiation is the radiolysis of water and the consequential elevation of the ECP caused by hydrogen peroxide production, which would lead to soluble CrO4 2− formation and depletion of Cr from the Cr-rich oxide phase. During reactor operation, the neutral pH would likely to be shifted to a more alkaline condition, which would push the equilibrium of Cr(III)/Cr(VI) towards a CrO4 2− favored situation, resulting in more Cr dissolution. Consequently, it would reduce the ability of FeCrAl alloy to form a protective spinel FeCr2 O4 oxide. This disruption of the formation of a protective barrier oxide would accelerate the overall corrosion

•

•

Oxide thickness and composition in the UR region after 24 hours of exposure is very close to the oxides found on FeCrAl alloys that have been tested in a long-term autoclave exposure study. However, the 24-hour exposure time is too short to establish a Cr-enriched region near the metal/oxide interface. The hydrogen peroxide formed by intense irradiation apparently has a dominant effect on the corrosion potential and, therefore, on the oxide composition. The absence of Cr in the oxide under irradiation is evidence of Cr(III) oxidation to Cr(VI), which is soluble as CrO4 2− . The formation of H2 O2 during irradiation of deaerated water has a more pronounced effect on corrosion potential, where the change is probably > 350 mV, and films that would have significant Cr3+ content will have little. Regardless of the presence of Ni, the oxide formed in both irradiation cases is consistent with the same elevation of ECP due to radiolysis products. The lack of difference between IR and RAR regions in the proton experiments indicates that the oxidation and oxide were controlled by the radiolysis products and not controlled by the point defect damage to the underlying metal.

Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Peng Wang: Writing - original draft, Writing - review & editing, Investigation, Data curation. Slavica Grdanovska: Data curation, Investigation, Writing - review & editing. David M. Bartels: Investigation, Writing - review & editing. Gary S. Was: Investigation, Writing - review & editing. Acknowledgments The authors would like to thank the staff of the Michigan Ion Beam Laboratory for their assistance in the conduct of experiments. This research was funded by the US Department of Energy’s Nuclear Energy University Program (grant number: DENE0 0 08272). This is manuscript #5290 of the Notre Dame Radiation laboratory. 11

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