Effects of Si, Mn on the corrosion behavior of ferritic–martensitic steels in supercritical water (SCW) environments

Journal Pre-proof Effects of Si, Mn on the Corrosion Behavior of Ferritic-Martensitic Steels in Supercritical Water (SCW) Environments Ziqiang Dong, Ming Li, Yashar Behnamian, Jing-Li Luo, Weixing Chen, Babak S. Amirkhiz, Pei Liu, Xin Pang, Jian Li, Wenyue Zheng, Dave Guzonas, Chenhui Xia

PII:

S0010-938X(19)32154-7

DOI:

https://doi.org/10.1016/j.corsci.2020.108432

Reference:

CS 108432

To appear in:

Corrosion Science

Received Date:

14 October 2019

Revised Date:

3 January 2020

Accepted Date:

3 January 2020

Please cite this article as: Dong Z, Li M, Behnamian Y, Luo J-Li, Chen W, Amirkhiz BS, Liu P, Pang X, Li J, Zheng W, Guzonas D, Xia C, Effects of Si, Mn on the Corrosion Behavior of Ferritic-Martensitic Steels in Supercritical Water (SCW) Environments, Corrosion Science (2020), doi: https://doi.org/10.1016/j.corsci.2020.108432

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Effects of Si, Mn on the Corrosion Behavior of Ferritic-Martensitic Steels in Supercritical Water (SCW) Environments

Ziqiang Dong1, Ming Li2, Yashar Behnamian2, Jing-Li Luo2, Weixing Chen2, Babak S. Amirkhiz3, Pei Liu3, Xin Pang3, Jian Li3, Wenyue Zheng4, Dave Guzonas5 and Chenhui Xia6

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Department of Chemical and Materials Engineering, University of Alberta, Edmonton, Canada MTL-NRCan, 183 Longwood Road South, Hamilton, Canada, Canada 4

Beijing University of Science and Technology, Beijing, China

School of Materials Science and Engineering, Shanghai University, Shanghai, China

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Highlights

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Effects of Si, Mn elements on corrosion of F/M steels in SCW were investigated. Si, Mn synergistically affected both the oxidation kinetics and oxide scales formed. Increasing Mn content reduced the corrosion resistance of F/M steels in SCW. Si/Mn ratio appears to be a useful indicator of long-term oxidation behavior of F/M steels. Water as a solvent must be considered when considering the stability of oxide films in the highdensity SCW.

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    

Consultant, Deep River, Canada

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2

Materials Genome Institute, Shanghai University, Shanghai, China

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1

Abstract:

Ferritic-Martensitic (F/M) steels with various amounts of alloying elements (Cr: 12 wt. %, Si: 0.6-2.2 wt. %, Mn: 0.6-1.8 wt. %) were exposed to Supercritical Water (SCW) environments at 500 °C and 25 MPa for up to 1000 hours. Gravimetric, Scanning Electron Microscope, Transmission Electron Microscopy (TEM), Energy-dispersive X-ray Spectroscopy (EDX) and X-ray Diffraction (XRD)

analyses were conducted to characterize the corrosion behavior of F/M steels in supercritical water. The results showed that the contents of the alloying elements (Si, Mn) affected both the oxidation kinetics and the oxidation scales formed on F/M steels. The F/M steel (Si: 0.6 wt. %, Mn: 1.8 wt.%) showed the highest corrosion rate in SCW with an intermittent oxide scale formed on the surface. Increasing the Si content in F/M steel improved the oxidation resistance and promoted the formation of a uniform oxide scale. The mechanism associated with the corrosion of these F/M steels under SCW conditions is

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discussed.

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Keywords: Corrosion, Supercritical Water, Ferritic-Martensitic Steels, Silicon, Manganese Introduction

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In recent years, the development of power plants (fossil-fired or nuclear) that operate at ultra-high

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temperatures and pressures has attracted increasing attention, aiming to improve the energy conversion efficiency and reduce pollutant emissions [1, 2]. Supercritical water (SCW), which refers to water above

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its thermodynamic critical point (374.1°C, 22.1 MPa), has been considered as a promising coolant for new next generation (Generation IV) nuclear power plants, with enhanced thermal efficiencies of 45-50

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% compared to current generation water-cooled reactors (WCRs) [3]. However, the proposed hightemperature and pressure operating conditions combined with irradiation present particular challenges to the materials used in a supercritical water-cooled reactor (SCWR) core since corrosion can be severe when metallic alloys [4] or ceramics [5, 6] are exposed to SCW, especially in the presence of high

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concentrations of dissolved oxygen (DO). Corrosion and stress corrosion cracking, which may lead to catastrophic failures in the power plant, can be difficult to predict in SCW environments.. A thorough understanding of the corrosion behavior of different candidate materials in SCW is essential for the development of SCWR concepts. Extensive investigations have evaluated a wide range of candidate alloys for use in SCWRs [4], including nickel-based alloys (e.g., Alloy 625) [7, 8], iron-based alloys (ferritic-martensitic steels [9-11], austenitic

stainless steels [12] and ODS steels [13, 14]), zirconium-based alloys [4, 15] and titanium-based alloys [4]. Various stainless steels (SS) have been selected as candidate in-core materials for the SCWR concepts proposed by Generation IV International Forum participants. Stainless steels exhibit high resistance against general corrosion because the high chromium content enables the formation of protective chromium oxide scales that reduce the corrosion rate. Extensive studies have been conducted to investigate the corrosion behavior of stainless steels in various SCW environments [4]. For example,

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Sun et al. investigated the oxidation of 316 SS in SCW with 2.0% H2O2 for up to 250 h with temperatures ranging from 350 °C to 500 °C [16, 17]. Their results showed that duplex oxide layers consisting of

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Fe2O3/(Fe,Cr)3O4/Cr2O3/Ni-enrichment/316SS (outer layer to inner layer) were formed after the SCW

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exposure. Chen et al. examined the corrosion behavior of the ferritic-martensitic steel T91 in SCW at 500 °C with two DO concentrations (25 ppb and 2 ppm) [18]. It was observed that a duplex oxide scale

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composed of an outer iron oxide layer and an inner spinel-dominated iron/chromium oxide layer formed after SCW exposure in the presence of 25 ppb DO. A tri-layer oxide scale consisting of an outer, porous,

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hematite layer followed by a magnetite layer and an inner iron/chromium spinel oxide layer was

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produced after exposure to 2 ppm DO.

Ferritic-martensitic steels (F/M) have attracted great attention because of their high creep resistance, high thermal conductivity, high radiation resistance and low thermal expansion coefficients [4]. However, studies have shown that the corrosion rates of these alloys are too high to allow their use at the high

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temperatures currently proposed for the various SCWR concepts. Although many studies have been conducted to investigate the corrosion mechanism of various F/M steels in SCW, few reported studies have examined the effects of alloying elements on the corrosion behavior of F/M steels in SCW. This study examined the corrosion behavior of F/M steels with various chemical compositions in SCW to elaborate the effects of the alloying elements Si and Mn on the corrosion.

2.

Experimental

F/M steels with the chemical compositions listed in Table 1 were evaluated in this study. The steels denoted as FM-(1~5) contain roughly the same amount of Cr (~11 wt.%) with variations in Si (0.6 - 2.2 wt.%) and Mn (0.6 -1.8 wt.%). Rectangular samples with a size of 20 mm×5mm×2mm were cut from as-received F/M steels and ground using silicon carbide papers from 240 to 1200 grit, followed by ultrasonic cleaning in ethanol for 30 minutes. SCW tests were conducted using tubular autoclaves made

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of 316L stainless steel tubing with an outer diameter of 9 mm and a wall thickness of 1.5 mm. The F/M steel samples were placed inside individual autoclaves filled with the amount of air-saturated deionized

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water (DO concentration of 8 ppm) required to achieve the test pressure (25 MPa) at the test temperature

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(500 °C). For each composition, three samples were tested for each duration and the average value of the weight gains were used in the analysis. The details of the SCW tests can be found in Ref. 6.

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The weights of the samples before and after SCW exposure were measured using a balance with five-

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decimal accuracy (0.00001g). The surface morphologies of the samples were characterized using a Zeiss EVO MA-15 Scanning Electron Microscopy (SEM) equipped with a Bruker Energy Dispersive X-ray

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Detector. X-ray diffraction (XRD) analysis was performed using a Rigaku Ultimate IV X-ray diffraction system. TEM characterization was performed using JEOL 2200FS TEM operated at 200kV equipped with energy dispersive spectroscopy. The TEM sample was prepared using the focused-ion-beam (FIB) lift-out technique (Hitachi NB5000 dual-beam FE/SEM). Results

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3.

3.1

Gravimetric measurements

Fig.1 (a) shows the oxidation kinetic plots obtained from the measured weight gains for all the samples exposed to SCW for times up to 1000 h. As shown in the figure, the weight gains for all the samples increased with increasing exposure time. After 1000 h exposure, Sample FM-5, containing the highest content of Mn and the lowest content of Si, showed the largest weight gain, while Sample FM-4,

which has the highest content of Si, exhibited the lowest weight gain. Comparing the weight gains of the three samples (FM-1, FM-2, and FM-5) which have the same amount of Si (around 0.7 wt. %), the weight gains generally increased with increasing Mn content. A similar comparison was made for the samples (FM-3, FM-4, and FM-5) which contain the same amount of Mn (around 1.8 wt.%), and the weight gain decreased with increasing the Si content. Suzuki [21] suggested that the oxidation behavior of steels containing Si and Mn was related to the Si/Mn ratio. In the present study, Sample

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FM-4, which has the highest Si/Mn ratio, showed the lowest weight gain among all the samples, and Sample FM-5, which has the lowest Si/Mn ratio, showed the highest. However, it is worth noting that

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although Sample FM-2 (Si/Mn: 0.6364) and FM-3 (Si/Mn:0.6111) have almost similar Si/Mn ratios,

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Sample FM-3 (Si, Mn: 1.1, 1.8) showed less weight gain than Sample FM-2 (Si, Mn: 0.7, 1.1), indicating that Si can reduce the detrimental effect of Mn on the oxidation of F/M steels in SCW when

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the absolute content of Si is increased. It should also be noted that the difference between the highest and lowest weight gains after 1000 h was only about 17 %, which would not be a sufficient

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improvement to allow use in an SCWR core.

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The weight gains of the samples were fit according to the generalized equations: ∆𝑊 = 𝑘𝑡 𝑛

and

−𝐸𝐴

(2)

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𝑘 = 𝑘0 𝑒 ( 𝑅𝑇 )

(1)

where k and k0 are rate constants, EA is the activation energy, R is the gas constant, T is the temperature, and n is an exponent that describes the time dependence of the weight change. The dependence plot of log(∆𝑊) on log(t) (Fig. 1(b)) shows that the first set of data points (100 h exposure) deviated from the linear curve fitted for the rest of the data sets, indicating that different kinetics existed at the beginning of SCW exposure. This difference could be related to the water chemistry change over the initial period of the SCW exposure because of the use of static autoclaves; the initial

oxidation rate may have been higher because the starting DO concentration was 8 ppm and oxygen was likely the primary oxidant. After the DO was largely consumed (within 250 h), water became the primary source of oxidizing species and the samples were exposed to an SCW environment with an extremely low DO concentration. During this initial stage of oxidation (denoted Stage I), both n and k showed a marked dependence on the concentration of Mn and on the Si/Mn ratio (Fig. 2a). From a Si/Mn ratio of 0.033 to 1.17, n decreased essentially linearly with increasing Si/Mn, while k decreased

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essentially linearly over the same range. At higher Si/Mn ratio, the value of k decreased abruptly, and

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the value of n increased. The values of n and k for the five samples are listed in Table 2.

The exponent n determined for the data points from 250 h to 1000 h (Stage II) ranged from 0.34 (FM-

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5) to 0.41 (FM-4). Fig. 2b shows that the values of n and k in Eq. 1 are a function of the Si/Mn ratio

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the value of k decreasing essentially linearly with increasing Si/Mn ratio, while the value of k showed a shallow minimum between Si/Mn rations of 0.6 to 1.2, The decreases in both n and of k results in a

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lower weight gain after 1000 h. Assuming no further changes in kinetics with exposure time, extrapolation of Eq. 1 to 25,000 h (a three-year fuel cycle) suggests that increasing the Si content of

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the steel from 0.6 to 2.2 wt.% would decrease the total corrosion at the end of component life by a factor of about 7.

In general the kinetics are similar to those reported in previous studies of F/M steels. Chen et al. [18]

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investigated the corrosion of T91 in SCW at 500 °C for up to 500 h. Their results showed that the oxidation kinetics of this alloy exposed to SCW with 2 ppm DO followed a parabolic behavior with an exponent n of 0.57, and a sample exposed to a much lower DO concentration of 25 ppb exhibited kinetics between parabolic and cubic behavior (n=0.40). Tan et al. investigated the corrosion behavior of 9-12% Cr F/M steels in SCW with a DO concentration of less than 25 ppb at temperatures ranging from 360 °C to 600 °C for up to 3000 h [19]. They found that parabolic oxidation kinetics existed for Alloy HCM12A (n=0.497) and Alloy NF616 (n=0.486).

3.2 Characterization of oxide scales Fig. 3 shows the surface morphologies of samples after 1000 h of SCW exposure. The surfaces of Samples FM-1, FM-2, and FM-4 appeared similar, being uniformly covered by polyhedral grains with a diameter of 3~5 µm. The surfaces of FM-1 and FM-4, whose Si/Mn ratios were greater than 1, were the most similar in appearance. The surface morphologies of Samples FM-3 and FM-5, which had the smallest Si/Mn ratios, were less uniform, with areas partially covered by polyhedral grains. Although the

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surface morphologies appeared different, the XRD patterns (not shown) obtained for all the samples were

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similar and only magnetite (Fe3O4) was indexed. Hematite is expected to be formed on steels exposed to SCW with DO concentrations greater than 2 ppm [20]. The observation of only magnetite on the surface

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is attributed to the presence of an SCW environment with an extremely low DO concentration after 250 h exposure. Li et al. investigated the evolution of the oxide scale of F/M steels in the early stage of SCW

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exposure with a DO concentration of 40 ppb [21]. They found that both magnetite and hematite were

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formed in the first 20 hours of exposure [21]. However, hematite disappeared completely, and only magnetite was detected when the exposure time extended to 40 hours.

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The cross-sections of the samples after SCW exposure are shown in Fig.4. Samples FM-1 and FM-4, with the smallest Si/Mn ratios, showed continuous uniform bi-layer oxide scales with a thickness of 14~16 μm. The oxide scales formed on Samples FM-2, FM-3 and FM-5 exhibited a less uniform structure compared to other samples, with more variation in the thickness, especially on Sample FM-5. A general

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increase of the oxide thickness was observed for Samples FM-3, FM-2 and FM-5 from about 20 μm (thickest part of Sample FM-3) to 33 μm (thickest part of Sample FM-5), which is consistent with the gravimetric measurements and the surface examination. A more uniform and protective oxide scale was observed on the surfaces of Samples FM-1 and FM-4 which showed lower weight gains than the other samples.

EDS analyses were conducted to characterize the oxide scales formed on Samples FM-4 (Si/Mn = 1.29) and FM-5 (Si/Mn = 0.33), which had the lowest and highest weight gains, respectively, after 1000 h exposure, and also had the highest and lowest Si/Mn ratios, respectively. The EDS mapping results for Sample FM-4 are shown in Fig. 5. The outer oxide layer was mainly composed of Fe and O with a trace amount of Mn, consistent with it being Fe3O4 with some Mn substitution in the lattice. The inner oxide layer consisted of an oxide (or oxides) containing Fe, Cr, Mn, and Si. It has been shown by others that

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the inner oxide layer on F/M steels consists mainly of an iron/chromium spinel [10, 11, 19, 22]. The nature of the incorporated Si could not be determined. The oxide scales formed on Sample FM-5

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exhibited a non-uniform structure with varying thickness as shown in Fig. 6. The oxide scale in the

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thicker part (Region A) consisted of three layers; a thick magnetite outer oxide layer, a layer enriched in Cr, Mn and Si and depleted in Fe that was roughly a mirror image in shape of the outer magnetite layer,

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and a thin layer highly enriched in Cr, somewhat enriched in Si, and strongly depleted in Fe that formed at the oxide-alloy interface. Less internal oxidation occurred in region B, which was enriched in Cr and

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Mn and depleted in Fe, indicating that a Cr-enriched compound, possibly chromia, was formed in this region that effectively inhibited the oxidation in this region. This is further supported by the observations

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in regions C and D; internal oxidation was still in progress in region C where no significant amount of Cr was detected but was arrested in region D because of the presence of a Cr-enriched layer. 3.3 In-depth examination of the oxide scale formed after SCW exposure

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As noted, the oxide scale formed on F/M steels generally exhibits a bi-layer structure consisting of an outer iron oxide layer and an inner spinel oxide layer which has been investigated extensively [18, 20, 21, 23]. As shown in Fig. 7, the typical bi-layer oxide scale formed on the surface of F/M steel in SCW consists of large angular grains composed of iron oxide with the longitudinal axis parallel to the oxide growth direction and which was formed by the outward migration of iron, whereas the inner oxide layer was composed of fine equiaxed Fe, Cr-spinel grains generated by the ingress of oxygen carriers [22].

The alloying elements Si and Mn were mainly distributed in the inner oxide layer as indicated by the EDS analysis, which shows enrichment of Si at the oxide/substrate interface. Beneath the inner oxide layer, there exists an internal oxidation transition layer at the alloy-oxide interface. With increasing Mn content, the oxide scale transitioned from a uniform and continuous structure (FM-4, FM-1) to a nonuniform and intermittent structure (FM-3, FM-2, FM-5) accompanied by an increase in weight gain. The intermittent oxide scale on Sample FM-5 (highest weight gain, lowest Si/Mn ratio) was different

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from that formed on the other samples. An in-depth characterization of the oxide scale formed on Sample

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FM-5 was performed using TEM/EDS. Fig. 8 shows the High Angle Annular Dark Field (HAADF) STEM image and EDS analyses on the oxide scale formed on Sample FM-5. As shown in Fig. 8, the

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structure of the oxide scale was clearly defined by two interfaces: Interface I (outer oxide layer/inner oxide layer) and Interface II (inner oxide layer/substrate). The HAADF STEM image provides a high-

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resolution image with the contrast reflecting the chemical composition change in different regions of the

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sample. As shown in Fig. 8 (a) and Fig. 8(d-1), the outer oxide layer consists of coarse oxide grains with the grain boundaries clearly visible in the HAADF image. The chemical composition of the outer oxide

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layer is quite uniform, iron oxide (magnetite) being the major compound formed due to the outward diffusion of Fe. The inner oxide layer exhibits a substantially different structure with a nonuniform chemical composition. The oxide grains in the inner oxide layer are indistinct in the HAADF image. As shown in the EDS mapping results (Fig. 7b), Fe was unevenly distributed in the inner oxide layer, being

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depleted in some areas where Cr was enriched, consistent with the EDS results shown in Fig. 8. The Cr was mainly enriched along the two interfaces (Interfaces I & II), indicating the formation of chromiumenriched oxides in those areas. As shown in Fig. 8(c), there is a noticeable Cr-enriched ‘island’ formed beneath Interface I. The center of the ‘island’ was enriched with Si while the contour was enriched with Cr. Conversely, Fe was depleted in this area. It was worth noting that EDS showed a sub-layer (band) structure in the inner oxide layer. This sub-layer structure is defined by two Cr-enriched “ribbons”. The

band structure reveals the dynamic diffusion process of cations in the inner oxide layer. In each sublayer, the Fe content increased gradually in the outward direction, while the Cr concentration decreased. The distributions of the alloying elements (Si, Mn) in the EDS mapping are difficult to discern due to their low concentrations. However, a narrow Mn-enriched region that matches the Cr-enriched region was observed along the interface region of oxide/substrate. To further identify the oxide composition variations within the oxide scale, quantitative EDS analyses were performed at the two interfaces.

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3.3.1 Interface I

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Interface I denotes the interface between the outer and inner oxide layers and is located at the original (unoxidized) steel surface. The elemental distribution images at Interface I are shown in Fig. 9. A thin

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Cr-enriched layer was observed at the original steel surface and at several regions underneath that

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surface. A very thin intermittent Si-enriched layer formed just beneath the Cr-enriched layer along the interface as well as in the Cr-enriched island area. Ellingham-type diagrams show that SiO2 is the most

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thermodynamically stable oxide of the elements present in the alloy, and would be expected to form first upon exposure of the sample surface to SCW, consistent with the observation of an Si layer below the

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Cr layer at Interface I. However, Ellingham diagrams do not take into account the role played by water as a solvent, rather than as a source of oxidant. Fournier and Potter [24] proposed an empirical equation for the solubility of quartz in water based on a reviewed of a large number of studies covering a wide range of temperatures and pressures (25 – 900 ◦C, pressures up to 10000 bar). Based on this study the

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solubility of SiO2 at 500 ◦C is on the order of ppm, suggesting that the silica layer forming on the steel surface undergo some dissolution, a process that may be exacerbated by the presence of a high concentration of DO at the start of the test. Work carried out by Wang et al. [25] on SiO2-doped Cr2O3 ceramics in high-temperature supercritical water provides support for this hypothesis. They reported that the dissolution of Cr decreased when SiO2 was added to Cr2O3 ceramics, but significant concentrations of Si were measured in the test solution. Chromium oxide would be expected to form next, although

again, at the high DO concentration present at the start of the test, some of the chromium may have dissolved into the SCW. Quantitative EDS analyses were performed at seven locations around the interface region as shown in Fig. 10. Table 3 lists the chemical compositions measured at the seven spots. Although Areas 1 and 2 were in the inner spinel region, both had high concentrations of Fe (41 at. %) with only minor amounts of Cr and Mn. The Si concentrations in the two areas differed significantly. The depletion of Cr in Areas

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1 and 2 is attributed to diffusion of Cr from the local region to the Cr-enriched region (Area 3). Area 3

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was the region where a short-circuit diffusion path (e.g. grain boundary) existed for the diffusion of oxidizing species; Cr preferentially diffused to those regions from adjacent grains and formed Cr-

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enriched oxides [21, 26]. A slight increase of Mn content (2.82 at. %) was also detected in Area 3 where Cr was enriched. A high Si content (12.09 at. %) was detected in Area 4 which was adjacent to the Cr-

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enriched region, and likely represents a portion of the silicon oxide that formed during the initial stage

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of the corrosion process. Area 5 was at the interface (outer oxide/inner oxide) region where chromium oxide was preferentially formed. No measurable Si was detected in the outer oxide layer (Areas 6 and

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7). Trace amounts of Cr and Mn were still present in the outer oxide layer near the interface region. The results of the detailed examination of Interface I can be summarized as follows: (1) The outer oxide layer was composed of magnetite (Fe3O4) doped with trace amounts of Cr and Mn; (2) Cr oxide was preferentially formed at regions on the steel surface and at grain boundaries where oxidizing species

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were present; (3) Si was only found in the inner oxide layer, silicon oxide forming adjacent to Cr-enriched oxides; (4) Mn was detected in both the inner and outer oxide layers, and was specifically enriched in regions where Cr was also enriched. 3.3.2 Interface II

Fig. 11 shows the elemental distribution at Interface II, which was the interface between the inner oxide layer and the alloy substrate. A distinct crevice was observed between inner oxide and alloy substrate, whose formation was attributed to the coalescence of the pores generated from the condensation of vacancies caused by the out-diffusion of cations [23]. As shown in the EDS mapping results, Cr was preferentially oxidized to chromium oxide along the crevice region on the alloy substrate side due to the presence of oxidizing species in the crevice region, as well as being preferentially oxidized on the grain

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boundaries. The signals for Si and Mn in the EDS mapping were too weak to be characterized.

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Quantitative analyses were conducted at different locations around the interface. Table 4 lists the chemical compositions at different spots near Interface II (Fig. 12). Area 1 was in the base metal and had

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the nominal alloy chemical composition. Area 2 was in the transition region where internal oxidation was in the process of intruding into the base metal. No measurable oxygen was detected in Area 2.

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Instead, a significant amount (93.04 at. %) of Fe was present in this spot. Area 3, which was located at

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the grain boundary area of the base metal, had a higher Cr content (21.75 at. %) than other areas indicating enrichment of Cr along grain boundaries. Quantitative analyses of Areas 2 and 3 indicated that

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there was diffusion of Cr from adjacent grains to the grain boundaries, which were preferentially oxidized due to the ingress of oxidizing species. Area 4 shows the chromium oxide precipitates formed beneath the crevice region. Similar to Area 2, Area 5 also had a high Fe content (66.79 at. %); this was the region where Fe was enriched due to the outward diffusion of Cr to the preferential oxidation region along the

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crevice. Area 6 and 7 were in the inner spinel layer region. Area 6 had higher Cr content (17.96 at. %) than Area 7, which was located closer to the oxide-SCW interface. These quantitative analyses revealed the inward oxidation progress due to the ingress of oxidizing species accompanied by the diffusion of cations and vacancies. Chromium, present in the alloy or diffusing from adjacent grains, was preferentially oxidized to chromium oxides near or at defects such as grain boundaries (Area 3), voids, or crevices (Area 4 and Area 6) due to the high affinity of Cr for oxygen. As

a result the metal grains were generally outlined by Cr-enriched oxides as shown in Fig. 6. Because of the diffusion of Cr from the metal grains (Area 2) to the grain boundaries (Area 3), the concentration of Fe was enriched inside the metal grains (Area 2). Oxygen could diffuse into the metal grains and react with Cr to form fine chromium oxide precipitates (A in Fig.11). The ingress of oxidizing species led to a sufficiently high oxygen partial pressure to oxidize Fe to fine equiaxed Fe3O4 and eventually formed the inner (Fe, Cr)3O4 spinel layer. No measurable Si was detected in Area 2, 4, and 5. A Si content of

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2.05 at. % was measured at the grain boundary area (Area 3) indicating the enrichment of Si at grain boundaries. With the exception of Area 2, all other areas contained trace amounts of Mn, especially in

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Area 3 (4.67 wt.%) and Area 4 (4.62 wt.%); Cr was also enriched in these regions suggesting the

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formation of Cr, Mn oxides.

Diffraction patterns (DP) of some selected areas in the sample were obtained. Fig. 13 shows the TEM

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images of the base metal, inner spinel layer, the outer magnetite layer, and the DP at different areas. The

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DP at Interface II shows nucleation of Fe3O4 at the interface of the base metal/inner oxide layer. The inner oxide layer consisted of nano-sized (Fe, Cr)3O4 spinel oxides as confirmed by the DP. The outer

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oxide layer was composed of larger magnetite oxides （Fe3O4）. Discussion

4.1

Thermodynamics of oxide formation

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Thermodynamic calculations can provide information on the stable phases expected to form on the surface of F/M steels exposed to an SCW environment. Fig. 14 shows the calculated isothermal phase diagram at a temperature of 500°C and 0.1 MPa. As few direct measurements of the species existing at equilibrium in metal alloy/ high density (e.g., 25 MPa) SCW exist, thermodynamic calculations relevant to high density SCW are based on data measured at lower temperatures and extrapolated to the conditions of interest using various theoretical models [27]. As a result significant uncertainties may exist in the

nature of the stable phases. Bearing these caveats in mind, the phase diagram calculated at 500°C and 0.1 MPa can provide insight regarding the chemical species expected to be stable at equilibrium at 500° C and 25 MPa. Fig. 12 (a) shows the ternary phase diagram of Fe-Cr-O. As shown in the diagram, the phases existing in the system include a BCC phase (solid solution of Fe and Cr), a M2O3 corundum phase, and spinel phases. Corresponding to the Fe-12wt. %Cr F/M steel, the sequence of phase formation with the

increase

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oxygen

is:

BCC;

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BCC+Corundum(Cr2O3);BCC+Corundum(Cr2O3)+Spinel((Fe,Cr)3O4);BCC+Spinel((Fe,Cr)3O4;BCC+

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Spinel((Fe,Cr)3O4+Spinel(Fe3O4);Corundum(M2O3)+Spinel(Fe3O4)+Spinel((Fe,Cr)3O4;Corundum(M2 O3). The phases calculated to be present were generally consistent with the oxides identified in the oxide

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scale on moving from the metal substrate to the outer surface with the increase of oxygen potential. As shown in Fig 9, Cr-enriched oxides (Area 3 and Area 4) were preferentially formed at the alloy/inner

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oxide layer interface adjacent to the substrate where the oxygen partial pressure was low. Along the

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cross-section from the substrate to the outer surface the spinel phases ((Fe，Cr)3O4) and (Fe3O4) existed. The ternary phase diagram of Cr-Mn-O at 500°C, 0.1 MPa was also calculated as shown in Fig. 13(b).

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With the increase of oxygen potential, the sequence of phases are: BCC+Spinel (MnCr2O4); BCC+Corundum(M2O3)+Spinel(MnCr2O4); Gas+Corundum(M2O3)+Mn2O3. The phase calculation suggested that the spinel phase MnCr2O4 can be formed even when only a low DO concentration was present, consistent with the observation in the quantitative analyses that Mn was enriched in those areas

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(Area 3 in Fig. 10, Areas 3 and 4 in Fig. 12) where a high amount of Cr was present. It was believed that the deterioration of the corrosion resistance of the F/M steels with increasing Mn content was related to the formation of Cr, Mn-containing oxide phases. 4.2

Oxidation of F/M steels in supercritical water

The bi-layer oxide scales observed for all the F/M steels studied were consistent with previously reported results on the oxidation of F/M steels in SCW [4, 11, 18, 19, 22] or high-temperature water vapor environments [28-30]. The outer oxide layer consisted of large angular grains predominantly composed of iron oxide, whereas the inner oxide layer was composed of spinel oxide phases. Beneath the inner oxide layer, a transition layer exists at the alloy-oxide interface (Fig.7). The outer iron oxide formed was attributed to the outward diffusion of Fe ions since iron shows a higher diffusion coefficient than that of

of

Cr ion in spinel phases [28], whereas the outward diffusion could also cause the inward transport of vacancies that condensed to form voids/crevices in the inner oxide scales. The inner oxide layer was

ro

formed by the inward penetration/diffusion of water molecules and oxygen produced by the dissociation

-p

of water via H2O=H2+1/2O2.

It is noteworthy that hematite was not observed on any of the samples in the current study, although the

re

initial DO concentration (8 ppm) was higher than the value (2 ppm) required to favor the formation of

lP

hematite over magnetite. The formation of magnetite instead of hematite in the current study is likely the result of the use of static tubular autoclaves in these tests. In order to estimate the SCW environment in

ur na

the sealed tubular autoclaves, calculations of reactant consumption were performed based on the geometry of the tubular autoclave and the experimental parameters, and the results are listed in Table 5. According to the calculations, if all of the DO (8 ppm) in the sealed system was completely consumed by the oxidation reactions, the resultant thickness of the oxide scale formed on the F/M steel surfaces

Jo

would only be 8-9 nm. The actual oxide scale formed on the F/M steels was much thicker (15 μm to 30 μm), which means that SCW was the primary source of oxidants during most of the test period. Although 8 ppm of DO was present at the beginning of the test, this oxygen was consumed during Stage I, and the DO concentration during Stage II was determined by the thermal decomposition of water. Li et al. also reported that hematite formed during the initial stage of SCW exposure gradually transformed to magnetite in an SCW environment with a low DO concentration (40 ppb), confirming that magnetite is

the favored product under this condition [21]. In a high-temperature and high-pressure hydrothermal environment, water dissociates via H2O=H2+1/2O2 to provide the oxygen for oxidation, and can directly react with the alloy, generating H2 as a by-product, resulting in accumulation of hydrogen at the oxide layer boundary [21]. It has been shown that even a small amount (1 ppm) of H2 could result in a significant decrease in the oxygen partial pressure (10-8 to 10-13) [29]. The presence of hydrogen may also influence diffusion processes occurring during the corrosion reaction, which could affect the

of

corrosion rate [31]. Hydrogen penetrates more easily into oxide layers and can influence the oxygen partial pressure at the metal-oxide interface. The potential effects of hydrogen are: (1) reducing the

ro

oxygen pressure at the metal-oxide interface which could cause preferential oxidation of Cr (Fig. 10,

-p

Area 4) in the bulk alloy; (2) facilitating iron transport, which has been observed when ferritic stainless steels were exposed to a moist hydrogen/ambient air environment [32]. Effects of Si and Mn on the corrosion of F/M steels in SCW

re

4.3

lP

By comparing all the samples evaluated, the general finding is that increasing the Si content increases the corrosion resistance of F/M steel in SCW and promote the formation of a more protective oxide scale

ur na

on the steel surface. Increasing Mn content caused the evolution of the oxide scale from a uniform/continuous structure to a non-uniform/discontinuous structure, accompanied by the deterioration of the corrosion resistance. The detailed characterization of Interfaces I and II provides some insights into the corrosion of F/M steels containing Si and Mn in SCW environments. The

Jo

interfacial reactions during the early stage of SCW exposure are reflected in the composition of Interface I, where thin layers of Si-enriched oxide and Cr-enriched oxide are observed. The Si/Mn ratio strongly influences the initial reaction rate through both the rate constant and the time exponent (Figure 2). These oxide layers slow but do not halt the oxidation process, hence, an outer magnetite layer doped with trace amounts of Mn due to the outer diffusion of Fe and Mn was formed. Meanwhile, oxygen carriers diffused

to the substrate preferentially along short-circuit paths (grain boundaries, voids, crevices), and Crenriched oxides were also formed in those areas. Although there has been no systematic study of the effect of Si on the oxidation behavior of steel under SCWR conditions, the influence of Si on oxidation processes occurring in low density SCW (superheated steam) and other oxidizing atmospheres (air, CO2) has been extensively studied [29,30]. The general finding is that adding a proper amount of Si increases the oxidation resistance of alloy due to the

of

formation of a protective silica layer; a vitreous silica layer formed at the oxide-alloy interface can act

ro

as a diffusion barrier during oxidation process [33-35]. The sequence of oxides observed at Interface I (Section 3.3.1) is consistent with this general finding. At the relatively high density of water present in

-p

the tests, this silica layer is not completely protective due to the significant solubility of silica under these conditions. The observation that Cr-enriched oxides were formed adjacent to the Si-enriched oxides

re

(Fig. 10) indicates that Si could promote the formation of Cr-enriched oxides. Swaminathan et al.

lP

investigated the selective oxidation of ternary Fe-Si-Cr alloys in an annealing environment at 820 °C in a N2-5% H2 gas atmosphere with various dew points [30] and found that increasing the Si content in the

ur na

bulk alloy promoted external Cr oxidation due to the formation of silica, which reduced the surface oxygen potential at the alloy-gas interface. As shown in Figure 5, the F/M steel sample with the higher Si content (2.2 wt. %) showed an obvious enrichment of Cr and Si in the inner oxide layer. Conversely, Sample FM-5, which has a lower Si content, showed an uneven distribution of Cr and Si in the inner

Jo

oxide layer.

The addition of Mn was found to accelerate the oxidation kinetics of F/M steels exposed to. Mn is a strong spinel-forming element that can promote the formation of MnCr2O4 spinel on Cr-containing alloys [36, 37]. The possible formation of MnCr2O4 may be indicated by the EDX results of Area 3 in Fig. 10, and Areas 3 and 4 in Fig.12. Gilewicz-Wolter et al. examined the diffusion rates of 51Cr, 54Mn and 59Fe in MnCr2O4 and FeCr2O4 and found that metal diffusion was higher in MnCr2O4 spinel [36]. The

presence of MnCr2O4 spinel in the inner oxide layer formed on the high-Mn steel may increase the diffusion of cations compared to that of FeCr2O4 spinel, preventing the formation of a protective chromia layer and promoting nodular growth of the oxide scale [36]. Additionally, as indicated by the EDX analyses, Mn has a high outward diffusion tendency which increases the transport of vacancies to the internal oxide layer. Condensation or segregation of these vacancies could create short-circuit paths

Conclusion

ro

5.

of

(voids, crevices, etc.) for the penetration of oxidants, accelerating internal corrosion.

The corrosion behavior of F/M steels with different Si (0.6 -2.2 wt.%) and Mn (0.6 -1.8 wt.%)

-p

concentrations was evaluated in SCW at 500 °C. The following conclusions were drawn from this

re

study:

(1) The kinetics of the oxidation of the F/M steels studied consisted of two stages; an initial period

lP

up to 100-250 h during which the kinetics varied between parabolic and linear depending on the alloy composition, and a second period from 250 h to the end of the test (1000 h) during which

ur na

the kinetics were between parabolic and cubic depending on the alloy composition. In both stages the rate constant and the time exponent were a function of the Si/Mn ratio. (2) Consistent with other studies of F/M steels in SCW, bi-layer structure oxide scales were formed

Jo

on steel surfaces with the outer oxide layer consisted of magnetite angular grains and the inner oxide layer composed of fine equiaxed spinel grains. Detailed characterization of the oxide scale formed on Sample FM-5 revealed that the outer magnetite layer was doped with trace amounts of Cr, Mn, and the inner layer structure was complex, consisting of a banded sub-layer structure caused by the diffusion of cations.

(3) The contents of alloying elements (Si, Mn) affected both the oxidation kinetics and oxide scales formed. Si improved the oxidation resistance of F/M steels in SCW, promoting the formation of a more uniform protective oxide layer. Increasing the Mn content reduced the corrosion resistance of F/M steels in SCW and resulted in non-uniform corrosion. The Si/Mn ratio appears to be a useful indicator of the long-term oxidation behaviour of these alloys. (4) The detailed characterization of the oxide scales revealed that both Si and Mn influence the oxide

of

microstructure. At high Si/Mn ratios (Si/Mn > 1) the oxide formed was relatively uniform and

ro

continuous, while at low Si/Mn ratios the oxide was less uniform, discontinuous and less protective. The role of Si is to form a somewhat protective silica layer during the initial stage of

-p

oxidation; the protectiveness of this layer is reduced due to the relatively high density of water which leads to some dissolution of the silica film. Manganese plays a role in the nature of the Cr-

re

bearing oxide formed, by promoting the formation MnCr2O4 in the inner oxide layer which may

lP

increase the diffusion of cations compared to that of FeCr2O4 spinel, increasing the corrosion rate. (5) The role of water as a solvent must be considered when considering the stability of oxide films

ur na

in the high-density SCW used as a coolant in SCWRs. Data availability

The raw/processed data required to reproduce these findings cannot be shared at this time due to technical

Jo

or time limitations.

Acknowledgments

The authors would like to acknowledge financial support from the NSERC/NRCan/AECL CRD program. The authors also appreciate Zhe Liu for conducting some of the experimental work.

References

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[1] R.B. Duffey, I.L. Pioro, Supercritical Water-Cooled Channel Nuclear Reactors: Review and Status, Japan, 2005. [2] J.M. Beér, High efficiency electric power generation: The environmental role, Progress in Energy and Combustion Science 33(2) (2007) 107-134. [3] A Technology Roadmap for Generation IV Nuclear Energy Systems; Generation IV International Forum, GIF-002-00, 2002 December. [4] G. Was, P. Ampornrat, G. Gupta, S. Teysseyre, E. West, T. Allen, K. Sridharan, L. Tan, Y. Chen, X. Ren, Corrosion and stress corrosion cracking in supercritical water, Journal of nuclear materials 371(13) (2007) 176-201. [5] Z. Dong, W. Chen, W. Zheng, D. Guzonas, Corrosion Behavior of Porous Chromium Carbide in Supercritical Water, Journal of Nuclear Materials 420(1-3) (2012) 320-327. [6] Z. Dong, W. Chen, W. Zheng, D. Guzonas, Corrosion behavior of chromium oxide based ceramics in supercritical water (SCW) environments, Corrosion Science 65(0) (2012) 461-471. [7] N. Boukis, W. Habicht, G. Franz, E. Dinjus, Behavior of Ni base alloy 625 in methanol supercritical water systems, Materials and Corrosion 54(5) (2003) 326-330. [8] J. Konys, A. Ruck, J. Novotny, J. Hausselt, Corrosion of Nickel Based Alloy G-30 Between 100 and 420 oC, pH Values From 2 to 6 and Pressures Ranging From 250-480 Bar Under Conditions of Supercritical Water Oxidation, CORROSION 2001 (2001). [9] P. Ampornrat, C.B. Bahn, G.S. Was, Corrosion and Stress Corrosion Cracking of Ferritic-martensitic Alloys in Supercritical Water, in: T.R. Allen, P.J. King, L. Nelson (Eds.) 12th Int. Conf. on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, 2005, pp. 13871396. [10] X. Ren, K. Sridharan, T. Allen, Corrosion of ferritic-martensitic steel HT9 in supercritical water, Journal of nuclear materials 358(2-3) (2006) 227-234. [11] K. Yin, S. Qiu, R. Tang, Q. Zhang, L. Zhang, Corrosion behavior of ferritic/martensitic steel P92 in supercritical water, The Journal of Supercritical Fluids 50(3) (2009) 235-239. [12] Y. Chen, K. Sridharan, T. Allen, Corrosion of Candidate Austenitic Stainless Steels for Supercritical Water Reactors, CORROSION 2007, NACE International, Nashville, Tennessee, USA, 2007, p. 14. [13] J. Isselin, R. Kasada, A. Kimura, Corrosion Behaviour of 16% Cr-4% Al and 16% Cr ODS Ferritic Steels under Different Metallurgical Conditions in a Supercritical Water Environment, Corrosion Science 52(10) (2010) 3266-3270. [14] J.H. Lee, R. Kasada, A. Kimura, T. Okuda, M. Inoue, S. Ukai, S. Ohnuki, T. Fujisawa, F. Abe, Influence of alloy composition and temperature on corrosion behavior of ODS ferritic steels, Journal of nuclear materials 417(1-3) (2011) 1225-1228. [15] Y. Jeong, J. Park, H. Kim, J. Busby, E. Gartner, M. Atzmon, G. Was, R. Comstock, M.G. da Silva, A. Motta, Corrosion of Zirconium-Based Fuel Cladding Alloys in Supercritical Water, in: T.R. Allen, P.J. King, L. Nelson (Eds.) 12th International conference on environmental Degradation of Materials in Nuclear Power System-Water reactors, TMS, 2005, pp. 1,369-1,377. [16] M. Sun, X. Wu, Z. Zhang, E.-H. Han, Oxidation of 316 stainless steel in supercritical water, Corrosion Science 51(5) (2009) 1069-1072. [17] M. Sun, X. Wu, E.H. Han, J. Rao, Microstructural characteristics of oxide scales grown on stainless steel exposed to supercritical water, Scripta Materialia (2009). [18] Y. Chen, K. Sridharan, T. Allen, Corrosion behavior of ferritic–martensitic steel T91 in supercritical water, Corrosion Science 48(9) (2006) 2843-2854.

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[19] L. Tan, X. Ren, T.R. Allen, Corrosion behavior of 9–12% Cr ferritic–martensitic steels in supercritical water, Corrosion Science 52(4) (2010) 1520-1528. [20] P. Ampornrat, G.S. Was, Oxidation of ferritic–martensitic alloys T91, HCM12A and HT-9 in supercritical water, Journal of Nuclear Materials 371(1–3) (2007) 1-17. [21] Y. Li, S. Wang, P. Sun, J. Yang, X. Tang, D. Xu, Y. Guo, J. Yang, D.D. Macdonald, Investigation on early formation and evolution of oxide scales on ferritic–martensitic steels in supercritical water, Corrosion Science 135 (2018) 136-146. [22] P. Ampornrat, C.B. Bahn, G.S. Was, Corrosion and stress corrosion cracking of ferritic-martensitic alloys in supercritical water, Proc. 12th Int. Conf. on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, held August (Warrendale, PA: The Minerals, Metals, and Materials Society [TMS], 2005), 2005. [23] J. Bischoff, A.T. Motta, C. Eichfeld, R.J. Comstock, G. Cao, T.R. Allen, Corrosion of ferritic– martensitic steels in steam and supercritical water, Journal of Nuclear Materials 441(1-3) (2013) 604611. [24] R.O. Fournier, R.W. Potter II, An equation correlating the solubility of quartz in water from 25 to 900 C at pressures up to 10,000 bars, Geochimica et Cosmochimica Acta 46(10) (1982) 1969-1973. [25] R. Wang, Z. Dong, J. Luo, W. Chen, W. Zheng, D. Guzonas, Stability of MgO-, CeO2-and SiO2doped Cr2O3 ceramics in high-temperature supercritical water, Corrosion Science 82 (2014) 339-346. [26] H. Hu, Z. Zhou, M. Li, L. Zhang, M. Wang, S. Li, C. Ge, Study of the corrosion behavior of a 18Croxide dispersion strengthened steel in supercritical water, Corrosion Science 65 (2012) 209-213. [27] L. Tan, Y. Yang, T. Allen, Oxidation behavior of iron-based alloy HCM12A exposed in supercritical water, Corrosion Science 48(10) (2006) 3123-3138. [28] A.S. Khanna, Introduction to high temperature oxidation and corrosion, ASM International (OH)2002. [29] J. Zurek, M. Michalik, F. Schmitz, T. Kern, L. Singheiser, W. Quadakkers, The Effect of WaterVapor Content and Gas Flow Rate on the Oxidation Mechanism of a 10Cr-Ferritic Steel in Ar-H2O Mixtures, Oxidation of metals 63(5-6) (2005) 5-6.

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face and Interface Analysis 40(3‐4) (2008) 268-272. [31] E. Essuman, G. Meier, J. Żurek, M. Hänsel, W. Quadakkers, The effect of water vapor on selective oxidation of Fe–Cr alloys, Oxidation of metals 69(3) (2008) 143-162. [32] Z. Yang, G. Xia, P. Singh, J.W. Stevenson, Effects of water vapor on oxidation behavior of ferritic stainless steels under solid oxide fuel cell interconnect exposure conditions, Solid State Ionics 176(17) (2005) 1495-1503. [33] M. Bennett, J. Desport, P. Labun, Analytical electron microscopy of a selective oxide scale formed on 20% Cr-25% Ni-Nb stainless steel, Oxidation of metals 22(5) (1984) 291-306. [34] A.M. Huntz, V. Bague, G. Beauplé, C. Haut, C. Sévérac, P. Lecour, X. Longaygue, F. Ropital, Effect of silicon on the oxidation resistance of 9% Cr steels, Applied Surface Science 207(1–4) (2003) 255-275. [35] A. Revesz, F. Fehlner, The role of noncrystalline films in the oxidation and corrosion of metals, Oxidation of Metals 15(3) (1981) 297-321. [36] A.L. Marasco, D.J. Young, The oxidation of Iron-Chromium-Manganese alloys at 900° C, Oxidation of Metals 36(1) (1991) 157-174. [37] H. Yearian, E. Randell, T. Longo, The structure of oxide scales on chromium steels, Corrosion 12(10) (1956) 55-65.

Tables and Figures

Table 1 Chemical composition of the F-M steels studied in this work Alloy Element (wt.%)

Name

Si

Mn

Fe

Ratio of Si/Mn

FM-1

11.1

0.7

0.6

balance

1.17

FM-2

11.2

0.7

1.1

balance

0.64

FM-3

11.6

1.1

1.8

balance

FM-4

11.6

2.2

1.7

balance

1.29

FM-5

11.7

0.6

1.8

balance

0.33

ro

-p

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of

Cr

0.61

Table 2. Fitting parameters for Eq. (1) used to fit the weight gain data for the five F-M steels studied during the initial (Stage I) and final (Stage II) stages of oxidation in current tests. Stage I

n K (mg cm-2h-n)

FM-1

FM-2

FM-3

FM-4

FM-5

FM-1

FM-2

FM-3

FM-4

FM-5

0.69

0.79

0.81

0.88

0.87

0.35

0.36

0.37

0.41

0.54

0.071

0.047

0.039

0.023

0.032

0.476

0.52

0.448

0.32

0.59

1

1

1

1

1

0.995

0.98

0.9772

0.991

0.999

Jo

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of

R2

Stage II

Table 3 Chemical composition of different areas at Interface I (Fig. 9).

Chemical Composition (at. %) Element Area 2

Area 3

Area 4

Area 5

Area 6

Area 7

Fe

41.67

40.79

4.92

12.79

31.46

44.39

44.08

Cr

3.03

0.57

29.19

3.9

6.02

0.45

0.27

Si

0.15

1.86

1.92

12.09

3.16

Mn

0.5

0.31

2.82

0.77

1.13

O

54.65

56.47

61.15

70.45

0.21

0.21

ro

55.44

58.23

-p re lP ur na Jo

of

Area 1

54.95

Table 4 Chemical composition of different areas at Interface II (Fig. 11). Chemical Composition (Atomic%) Area 1

Area 2

Area 3

Area 4

Area 5

Area 6

Area 7

Fe

79.59

93.04

7.18

1.67

66.79

11.63

26.53

Cr

12.21

1.2

21.75

49.2

8.34

17.96

10.22

Si

0.92

2.05

1.04

0.51

Mn

2.49

4.67

4.62

1.23

62.98

16.93

23.32

2.78

1.66

65.2

59.33

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Element

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Table 5 The calculation of the consumption of oxygen or water to generate oxide layer on the surface

Density

Thickness of the Water amount Percentage of water oxide scale formed needed to generate consumed to generate (g/cm3) on the sample by an 18 µm-thick a 18 µm-thick oxide the oxygen (8 ppm) oxide layer on the layer (assumed tube dissolved in water surface of the length is 20 cm) (%) (µm) sample

-p

Molecular mass (g/mol)

re

Compound

ro

of an F/M steel sample in the tubular autoclave.

231.53

5.1

Fe2O3

159.69

5.24

Cr2O3

151.99

5.21

223.84

4.8

Jo

FeCr2O4

0.0096

14.27

2.82

0.0086

15.94

3.142

0.0082

16.66

3.29

0.0099

13.90

2.74

ur na

Fe3O4

lP

(µL)

of ro -p re lP ur na Jo

Fig. 1 (a) weight changes of samples during the SCW exposure up to 1000 h at low oxygen level; (b) Logarithm plot of weight change against exposure time.

of ro -p re lP ur na Jo Fig. 2 Exponent n and rate constant k in Figure 1 for Stage 1 and Stage 2 of oxide growth.

of ro

Jo

ur na

lP

re

-p

Fig. 3 Surface morphologies of samples after SCW exposure at 500°and 25 MPa for 1000 h.

of ro

Jo

ur na

lP

re

-p

Fig. 4 Cross-sections of oxides formed on specimens after SCW exposure at 500°and 25 MPa for 1000 h.

of ro

Jo

ur na

lP

re

-p

Fig. 5 EDS mapping results of the oxide scale formed on Sample FM-4 after 1000 h exposure to SCW at 500°and 25 MPa.

of ro

Jo

ur na

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re

-p

Fig. 6 EDS mapping results of the oxide formed on Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa.

of ro -p

re

Fig. 7 EDS line-scan results of the oxide scale formed on Sample FM-1 after 1000 h exposure to SCW at 500°and 25 MPa.

Jo

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.

of ro -p re

Jo

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Fig. 8 Characterization of the oxide scales formed on Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa. (a) HAADF STEM image; EDS element mapping results of (b) Fe; (c) Cr; (e) Mn; (f) Si; HAADF STEM image of Interface I (d-1) and Interface II (d-2).

of ro -p re lP ur na

Jo

Fig. 9 Element mapping analyses at Interface I on Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa.

of ro -p

Jo

ur na

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re

Fig. 10 Quantitative analyses of different spot areas at Interface I on Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa.

of

A

B C

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re

C

-p

B

ro

A

Fig. 11 Element mapping analyses at Interface II on Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa; A: Cr-enriched spinel oxides formed inside metal grain; B: Cr-enriched spinel oxides formed along the grain boundary; C: Cr-enriched spinel oxides precipitates formed near the fissure regions.

Fig. 12 Quantitative analyses of different spot areas at Interface II on Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa.

Jo

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Fig. 13 TEM image of (a) base metal/spinel; (b) spinel phase; (c) spinel/magnetite; and diffraction patterns obtained for the selected area as indicated by the arrow: (d) Fe3O4 phase; (e) (Fe, Cr)3O4 spinel; (f) Fe3O4. Images obtained from Sample FM-5 after 1000 h exposure to SCW at 500°and 25 MPa.

（a）

(b)

Fig. 14 Calculated isothermal phase diagram at 500°C and 0.1 MPa: (a)1-BCC+Corundum(M2O3); 2BCC+Corundum(M2O3)+Spinel((Fe,Cr)3O4); 3-BCC+Spinel((Fe,Cr)3O4); 4-BCC+ Spinel(Fe3O4)+Spinel((Fe,Cr)3O4; 5-Corundum(M2O3)+ Spinel(Fe3O4)+Spinel((Fe,Cr)3O4; 6Gas+Corundum(M2O3);

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(b)Cr-Mn-O: 1-BCC+Spinel (MnCr2O4); 2-BCC+Corundum(M2O3)+Spinel (MnCr2O4); 3Gas+Corundum(M2O3)+Mn2O3