Effect of surface machining on the corrosion behaviour of 316 austenitic stainless steel in simulated PWR water

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Effect of surface machining on the corrosion behaviour of 316 austenitic stainless steel in simulated PWR water Siyang Wanga, Yujin Hua, Kewei Fangb, Wenqian Zhanga, Xuelin Wanga,

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a State Key Laboratory of Digital Manufacturing Equipment and Technology, School of Mechanical Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China b Suzhou Nuclear Power Research Institute, Suzhou 215004, China

A R T I C L E I N F O

A B S T R A C T

Keywords: A. Stainless steel B. TEM B. SEM B. XPS B. XRD C. Oxidation

The corrosion behaviour of 316 austenitic stainless steel with different surface machining processes in simulated pressure water reactor primary environment was investigated. The oxide films related to three surface conditions were characterized by using SEM, TEM and XPS. The results showed that the machined surface conditions affected the thickness and elemental distribution of the oxide films composed of a double-layer structure. The more heavily cold worked layer results in less chromium enrichment in the inner film, and more nickel enrichment at the oxide/metal interface. Meanwhile, unduly chromium depletion beneath the oxide film was observed on the machined surfaces.

1. Introduction Austenitic stainless steels are widely used in nuclear power plants to manufacture critical components owing to their corrosion resistance at high temperature. The corrosion resistance of austenitic stainless steels is mainly attributed to the oxide film formed on the alloy surface. Recently, the oxide film of austenitic stainless steels in high temperature water has been examined by numerous researchers [1–8]. The results reveal that oxide film formed on austenitic stainless steel has a duplex-layer structure. The outer layer of the film is mainly composed of large oxide particles which are magnetite (Fe3O4) or iron–nickel spinel oxide, and the inner layer is made up of nanocrystalline spinel which is rich in chromium. In particular, the chromium-rich layer is thought to play an important role in corrosion resistance. Despite the fact that corrosion behaviour of the material mainly depends on the material itself and corrosion environment, actually the corrosion of the material may be affected by the factor which is nonchemical in nature, such as the microstructure, the surface topography, and the residual stress. Ralston and Birbilis [9] reviewed the influence of grain size on corrosion, and suggested that the effect is varied depending on processing and environment even for the same material. Schion and Kenny [10] reported that pitting corrosion resistance and intergranular corrosion resistance of 304 austenitic stainless steel were improved by grain refining despite that general corrosion resistance is impaired by it. Lv and Luo [11] stated that the grain refinement, obtained by rolling mill and anneal, increased the concentration of

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chromium in the oxide film of 304 stainless steel after polarization at 0.2 VSCE for 12 h. As a result, the corrosion resistance of stainless steel was improved by the grain refinement. Li et al. [12] reported that for 316L austenitic stainless steel with nanometre-sized grains fabricated by means of dynamic plastic deformation, the spontaneous passivation ability and growth rate of the oxide film formed after polarization at 0.5 VSCE for 30 min were improved compared with the normal 316L austenitic stainless steel. Additionally, surface topography is another well-known factor affecting surface corrosion. Burstein et al. [13,14] studied the effect of roughness on the pitting of stainless steel in chloride solution. They reported that a smoother surface reduced the incidence of metastable pitting and prevented the propagation of nucleated sites. Leban et al. [15] found that the pitting potential of 304 stainless steel depended strongly upon the surface roughness in NaCl aqueous solution, but this was not true for the test in simulated urban rain. Li et al. [16,17] suggested that the “peak” of the fluctuating surface was more likely to corrode than the “valley” on a rough surface. Enlargement of surface roughness expedited the surface corrosion. Because the physical and chemical properties of oxide films play important roles in stress corrosion cracking (SCC) which is caused by the synergy of tensile stresses and a corrosive environment, it will be necessary to better understand the impact of the changes of surface and near-surface conditions on the microstructure of oxide film and the implications for material performance. To achieve better mechanical performance, almost all the components employed in nuclear power plants are produced by surface

Corresponding author. E-mail address: [email protected] (X. Wang).

http://dx.doi.org/10.1016/j.corsci.2017.06.019 Received 1 January 2017; Received in revised form 22 June 2017; Accepted 25 June 2017 0010-938X/ © 2017 Elsevier Ltd. All rights reserved.

Please cite this article as: Wang, S., Corrosion Science (2017), http://dx.doi.org/10.1016/j.corsci.2017.06.019

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machining (surface cold working) such as milling, grinding and mechanical polishing. In addition to the change of surface topography or roughness, surface machining generates a cold worked layer with a thickness ranging from about 100 μm to 300 μm. This cold worked layer has different mechanical and metallurgical properties from the original material, including the size of grain, dislocation density, and residual stress, etc. The changes of surface properties due to surface machining were found to have great influence on the metal’s corrosion behaviour in high temperature water. Guo et al. [18] studied the effect of surface states on the oxidation behaviour of 316LN stainless steel in high temperature pressurized water, and reported that the oxidation rate was reduced by electropolishing due to the higher content of chromium in oxide film. Han et al. [19] contrasted the corrosion of 316L stainless steel between electro-polished surface (EPS) and colloidal silica slurry polished surface (CPS), and stated that corrosion rate was higher for the EPS than the CPS. Ming et al. [20] reported that the oxide decreased with increasing of dislocations and subgrain boundaries in the cold work layer for 308L welded metal with different surface states in simulated pressure water reactor (PWR). Han et al. [21] suggested that electropolishing could improve the oxidation resistance of 316L stainless steel in simulated PWR water especially during the early oxidation stage. Warzee et al. [22] reported that the surface treatments accompanied by cold-working, milling and grinding, could reduce the corrosion rate of stainless steel in superheated steam at temperatures between 400 °C and 600 °C. Ziemniak et al. [3,23] compared the corrosion kinetics of 304 stainless steel in mildly alkaline, hydrogenated water at 260 °C between electro-polished surface and machined one. They concluded that the corrosion rate of electro-polished austenitic stainless steel was far below the machined one. Cisse et al. [8] studied the effect of surface preparation on the corrosion of austenitic stainless steel 304L in high temperature steam and simulated primary water of PWR, and reported that polishing promoted the development of oxide film. Ghosh et al. [4] studied the oxide films formed on machined and ground stainless steel in high purity water at 300 °C and pressure of 10 MPa. They suggested that the oxide films grown on machined and ground stainless steel had higher specific resistivity than solution annealed one. These previous studies highlighted the impact of surface machining on the properties of oxide films formed on austenitic stainless steel with different surface preparations. However, changes of surface properties introduced by surface machining are complicated, and their effects on the corrosion behaviour have not been under thorough investigation. More works are needed to establish a correlation between the changes of surface properties induced by machining and the microstructural characterization of the oxide film. Furthermore, properties of corroded subsurface (or the matrix beneath oxide film) also directly affect the exfoliation behaviour of oxide film and SCC crack initiation and propagation. Therefore, it is necessary to evaluate the effect of surface machining on the oxide/metal interface and subsurface. This study aims to explore the effect of the surface property changes caused by surface machining on the corrosion behaviour of 316 austenitic stainless steel in simulated PWR water. Three kinds of samples including electro-polished, ground and milled surfaces were considered. The surface morphologies and microstructures of the cold worked layer were characterized by surface roughness, micro-hardness and grain size. An oxidation test was carried out under simulated PWR water, the microstructure of the oxide film was characterized by field emission scanning electron microscope (SEM), field emission transmission electron microscope (TEM) and X-ray photoelectron spectroscope (XPS).

Table 1 Chemical composition of the AISI316 stainless steel (wt.%). Fe

Cr

Ni

Mo

Mn

Si

C

S

P

67.75

16.92

11.61

2.22

1.06

0.35

0.074

0.01

≤0.002

30 mm × 10mm × 3.5 mm using electrical discharge machining. Then they were solution-treated at 1040 °C for 30 min, followed by airquenching before being subjected to different surface finishing. Three types of surfaces were introduced in this study as (a) electropolished surface (denoted by “P”), (b) ground surface (denoted by “G”), and (c) milled surface (denoted by “M”). The first type of surface was ground with 400, 800, 1200–2000 grit waterproof abrasive papers, and then electro-polished by using 20 vol.% perchloric acid (HClO4) + 80 vol.% acetic acid (CH3COOH) to remove the Beilby layer, which was produced by the grinding process. The electro-polish voltage was kept at about 35 V for 20 s. The removed material thickness was about 5–10 μm. The second surface was directly ground by a grinding machine. The spindle speed was 1400 r/min, the cutting depth was 0.01 mm, and the diameter of grinding wheel is 250 mm. The last one was machined by a milling machine with a carbide tool. The spindle speed was 600 r/min, the feed rate was 30 mm/min, and the cutting depth was 0.1 mm. Before immersion test, all samples were washed by an ultrasonic cleaner with ethanol and deionized water for three times. The surface roughness of the samples was measured by using a comprehensive measurement system for surface profiler (Form Talysurf PG 1830). Sampling length was 1 cm. The measuring direction was along the machining direction. Profile arithmetic average error Ra was chosen to represent the surface roughness. The surface Vickers micro-hardness (HV) of the samples was measured by using a micro-hardness tester equipped with a diamond pyramidal indenter. A load of 1 N was applied for 15 s. Every sample was measured in different regions for four times. The diagonal of the rhombic indentation was about 20 μm. So the calculated depth of indentation was about 3 μm which was sufficiently small to ensure that only the machining affected layer was measured. The grain size of the surface cold worked layers was analysed by using an X-ray diffractometer. And the x-ray radiation source was Cu Kα with the applied voltage of 40 kV and the current of 30 mA. The scan ranged from 38° to 100° with a speed of 2 deg/min. Since the grain refinement and microstrain lead to line broadening of Bragg diffraction peaks, the average grain size can be calculated by using the Williamson–Hall method below [24].

β cos θ 1 2 sin θ = + 2ε⋅ λ D λ Where β denotes the integral breadth of the measured peak, θ the diffraction angle, λ the wavelength, D the grain sizes, ε the microstrain, and 2sin θ/λ the reciprocal lattice spacing. If β cos θ/λ is regarded as the Y axis, and 2sin θ/λ is regarded as X axis, then the reciprocal of grain size 1/D is the Y-intercept, and the microstrain ε is the slope of the linear equation. As there are multi Bragg diffraction peaks, each Bragg diffraction peak can be fitted to obtain an (x, y) coordinate pair. In this study, five Bragg diffraction peaks were selected to take the least squares fitting. The grain size of bulk material was observed under a light microscope. Firstly, the gridding with a spacing of 25 μm was drawn on the image of the bulk material. Then the grain sizes were measured on every gridding line in depth direction and surface direction, respectively. The surface direction is perpendicular to the depth direction as displayed in Fig. 3. At last, a statistical method was used to calculate the grain size in each direction. Moreover, the grain size of the cold worked layer was also measured on the TEM images with the same method, but the grid spacing of 25 nm was used. The oxidation test was carried out in a 3 L autoclave. 1.75 L

2. Methods The samples used in this study were cut from a 316 austenitic stainless steel tube. The chemical composition is given in Table 1. The rough samples were fabricated with dimensions 2

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Fig. 1. Schematic of TEM sample preparation.

deionized water, with 2.2 ppm (by mass) Li+ (LiOH·H2O) and 1000 ppm (by mass) B3+ (H3BO3) added, were poured into the autoclave to simulate the PWR water environment. 99.999% high purity argon gas was introduced into the solution to remove oxygen until the oxygen concentration of the solution below 5 ppb (by mass). No hydrogen was added in this study. The test temperature and pressure were maintained at 300 °C and 10 MPa, respectively. The heating speed was about 100 °C/h. The pressure control was conducted using a personal computer-based monitoring unit. After oxidation for 120 h, the autoclave was left to cool naturally at room temperature. At last, all samples were taken out and washed by deionized water carefully. In order to observe the oxide film microstructure on different surfaces, SEM (FEI Nova Nano SEM 450) analysis was performed in secondary electron (SE) mode with an accelerating voltage of 10 kV. For cross-sectional TEM analysis, the samples of the oxide film were prepared by using SEM/FIB dual-beam system (FEI Quanta 3D FEG). The schematic of TEM sample preparation was illustrated in Fig. 1. An area of interest was firstly located and a layer of platinum was deposited on the selected surface region to protect the oxide film from Ga ion sputtering (Fig. 1a). The deposited area was with a size of 6 μm × 2 μm. In this study, the areas of interest were chosen along machining marks. Then, triple prism grooves with a depth of about 4 μm were made on both sides of the Pt deposited area by using the focused ion beam (FIB) with Ga ion sputtering, and the lamella with platinum clad layer was left (Fig. 1b). Subsequently, the lamella was cut free from the matrix (Fig. 1c). It should note that the lamella was Pt welded to probe before it was completely cut off. At last, the lamella was Pt welded on a TEM sample grid and thinned from 2 μm to 100 nm by Ga ion sputtering (Fig. 1d). A field emission transmission electron microscope (FEI Tecnai G2 F30) with an energy dispersive X-ray spectroscopy (EDX) was used to observe the cross-sectional morphology of the oxide film and analyse its elemental composition with an accelerating voltage of 300 kV. The observation of morphology was in bright field mode, and the analysis of

the elemental composition was in STEM mode. During the TEM analysis, electron diffraction was performed to analyse the oxide particle of the outer oxide film. EDX line scan analysis was performed to clarify the elemental distribution of the cross-sectional samples. Additionally, EDX point analysis was used to quantitatively compare the elemental composition differences among the oxide films formed in three types of surfaces. Three EDX points were arranged in the outer layer, inner layer, and oxide/metal interface respectively. The points beneath the oxide film spaced a few tens of nanometres along the depth direction, as displayed in Figs. 7–9, were also analysed. The thickness of inner oxide layer was measured on TEM images. As the thickness of inner oxide layer was somewhat varied in different region, a statistic calculation was used to assess the thickness of inner oxide layer. 40 regions of each sample were chosen to measure the thickness of inner oxide film. Every measuring region was about 30 nm apart from each other. The valence states of elements in the oxide film were measured by XPS (AXIS-ULTRA DLD-600W). Photoelectron emission was excited by monochromatic Al Kα radiation (E = 1486.6 eV). XPS analysis was undertaken for each surface at the unsputtered oxide film and sputtered oxide film respectively. The sputtering time of 300 s was adopted to preferably present the oxidation states of elements in the inner oxide film. The argon ion sputtering with an ion energy of 500 eV was conducted over an area of 2 mm × 2 mm for 300 s. Sputtering rate was about 0.01 nm/s with reference to Ta2O5 layer. The survey scan from 0 to 1200 eV binding energy was performed on each sample. Then the O1s, Cr2p, Fe2p, Ni2p, Mo3d and C1 s detailed spectra were obtained in a high resolution of 0.48 eV/(Ag 3d5/2)@400 kcps. The analysis area was about 300 × 700 μm. The energy scale was calibrated by placing the C1s peak at 285 eV.

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Fig. 2. Surface micro-hardness and surface roughness of three kinds of surface before oxidation.

3. Results

respectively. From these data, it is apparent that the cold working of milled surface was more heavily than the ground one.

3.1. Characterization of different surfaces before oxidation 3.2. Cross-sectional morphologies of cold worked layer Fig. 2 shows the measurement results of surface roughness and surface Vickers micro-hardness of three kinds of surface before oxidation. The surface roughness increased gradually from electro-polished surface, ground surface to milled surface. The surface Vickers microhardness of electro-polished surface was the lowest and that of milled surface was the highest. For ground and milled samples, the mean grain size measured by XRD was 107.7( ± 19.1) and 90.1( ± 13.4) nm,

Fig. 3a is the cross-sectional metallographic observation of bulk material under a light microscope, showing the typical austenitic structure. The mean grain size along the depth direction is 64.87 μm while the mean grain size along the surface direction is 63.6 μm. The total average grain size is 64.28 μm. The ratio of grain sizes in two directions is 0.99. It is clear that the austenite grain is almost equiaxed.

Fig. 3. Cross-sectional morphology observation of matrix. (a) Image of bulk material of electro-polished sample under a light microscope. (b) TEM image of cross-sectional morphology of matrix of electro-polished sample. (c) Diffraction pattern showing that the crystal is austenite (zone axis [−1 1 1]). Frequency distribution histogram image of grain size refers to: (d) grain size along surface direction, (e) grain size along depth direction, and (f) total grain size.

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Fig. 4. Cross-sectional morphology of cold worked layer of ground sample. (a) Image of cold worked layer of ground sample under a light microscope. (b) TEM image of cross-sectional morphology of cold worked layer of ground sample. (c) Diffraction pattern showing that the polycrystal is austenite. Frequency distribution histogram image of grain size refers to: (d) grain size along machining direction, (e) grain size along depth direction, and (f) total grain size.

3.3. Surface morphologies of oxide film

Fig. 3b is a TEM image of the matrix of electro-polished sample, and the matrix is a single crystal of austenite according to the electronic diffraction pattern (Fig. 3c). Fig. 4a shows the cross-sectional metallographic observation of the cold worked layer as a result of ground machining under a light microscope. There is a layer with a thickness of less than 10 μm on the surface which is different from the matrix. Numerous studies reported that this layer has the nanocrystalline structure introduced by surface cold working [25,26]. Fig. 4b is the TEM image of cross-sectional morphology about 20 nm below the ground surface, which shows the presence of nano-sized grains. The original grain was fractured due to grinding and the grain boundary was quite ambiguous. Meanwhile, the dislocation density was relatively high in the subsurface compared with the bulk material. The electronic diffraction pattern shows the cold worked layer is polycrystalline austenite. Statistically, the mean grain sizes in the machining and depth direction are 87.83 nm and 64.31 nm, respectively. The total average grain size is 75.38 nm. The grain size along machining direction is larger than that of the depth direction. The ratio of grain sizes in two directions is 1.37. Fig. 5a is the cross-sectional metallographic observation of the cold worked layer of milled sample under a light microscope. A layer with nano-sized grains about 10 μm thick was also observed. Fig. 5b is the TEM image of cross-sectional morphology of milled sample (about 20 nm below the oxide film). The grain was also fractured and the grain boundary was ambiguous. A large number of dislocations were observed. The electronic diffraction pattern shows the cold worked layer is polycrystalline austenite. The mean grain size along the machining direction is 91.24 nm, while the mean grain size along the depth direction is only 61.09 nm. The total average grain size is 72.54 nm which is a bit smaller than the ground one. The ratio of grain sizes in two directions is 1.49. The elongation of grain along the machining direction is more severe than that of ground one.

Fig. 6 shows the SEM images of typical surface morphologies from three different surfaces after exposure to simulated PWR high temperature water for 120 h. At low magnification, the structure of oxide products was not observable, but machining marks were still clear for both ground and milled surfaces (Fig. 6d, g). The electro-polished surface (Fig. 6a) was relatively smooth. At high magnifications, the electro-polished surface was sparsely covered with oxide particles (Fig. 6b, c) ranging from 150 to 165 nm, and the size of the oxide particles was quite uniform. The ground surface was also covered with oxide particles which were denser than that of the electro-polished one (Fig. 6e, f), and the size of the oxide particles ranged from 50 to 280 nm. For milled surface, it was densely covered with the oxide particles like ground one (Fig. 6h, i), and the uniformity of the oxide particle size was the poorest among three samples with a size ranged from 50 to 400 nm. 3.4. Cross-sectional morphologies of oxide film Fig. 7 shows a cross-sectional TEM image of the oxide film grown on the electro-polished surface. The oxide film consisted of a double-layer structure. The outer layer was composed of single oxide particles. The electron diffraction pattern of oxide particle appeared as a spot image, which confirmed that the particle was a single crystal. Meanwhile, the electron diffraction pattern demonstrated that the crystal was magnetite (Fe3O4, Orthorhombic). It should be noted that the electron diffraction was conducted at several oxide particles to avoid fortuity of the result. The inner layer was composed of nanocrystalline which was compact, and the thickness of inner oxide layer was quite uniform with approximately 23 nm. Fig. 8 is the cross-sectional TEM image of oxide film formed on the 5

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Fig. 5. Cross-sectional morphology of cold worked layer of milled sample. (a) Image of cold worked layer of milled sample under a light microscope. (b) TEM image of cross-sectional morphology of cold worked layer of milled sample. (c) Diffraction pattern showing that the polycrystal is austenite. Frequency distribution histogram image of grain size refers to: (d) grain size along machining direction, (e) grain size along depth direction, and (f) total grain size.

Figs. 11–13). As the materials of Pt layer, outer oxide film, inner oxide film and matrix are different, the thickness reduction of sample at each layer is different when being thinned by focused ion beam (FIB) with Ga ion sputtering. The count of STEM HAADF electron signal at one point has a positive correlation with the thickness of sample at this point, so the STEM HAADF electron profile can be used to differentiate different layers. Compared the results of EDX line scan, the elemental distributions of three kinds of oxide films were similar: (a) the outer oxide film was principally composed of iron oxide, which agrees with electron diffraction pattern analysis, (b) The inner oxide film was rich in chromium. The concentration of molybdenum in the inner oxide film was fairly high in comparison to other layers, and (c) nickel enrichment was observed across the oxide/metal interface. The thickness of nickel enrichment zone in the electro-polished sample was less than 10 nm, and that of ground sample was about 15 nm. The nickel enrichment zone of the milled sample was the biggest among three kinds of samples, up to 30 nm in thickness. For purpose of accurately comparing the elemental composition differences among three kinds of oxide films, EDX point analysis was performed further. It should be pointed out that as the TEM sample would inevitably contact oxygen in the atmosphere during sample preparation, the content of oxygen was not included, and only the metallic elements of iron, chromium, nickel and molybdenum were included (the weight percent of this four metallic elements is above 98.5% in 316 austenite stainless steel). The measured weight percent of every metallic element was transformed into the relative weight percent, which was calculated by normalized the concentration of each metallic element to the total concentration of the four metallic elements at the measurement point. Fig. 14 shows the relative weight percent profiles of four metallic elements across the oxide film. The oxide/metal interface was

ground surface. The oxide film cross-sectional morphology was also a double-layer structure: an outer layer composed of single oxide particles and an inner layer composed of compact nanocrystalline. The electron diffraction pattern of the outer layer also demonstrated that the crystal was magnetite (Fe3O4, Orthorhombic). The thickness uniformity of the inner layer was relatively poor compared with electropolished one. Fig. 9 shows the cross-sectional TEM image of oxide film formed on the milled surface. Duplex-layer structure was still observed. The singlecrystal oxide particles were magnetite (Fe3O4, Orthorhombic), and the thickness of inner oxide layer was rather nonuniform. The frequency distribution histograms of the inner oxide film thickness grown on different surfaces are shown in Fig. 10. The mean thicknesses of inner oxide layers on the electro-polished surface, ground surface and milled surface were 22.97 nm, 24.34 nm and 26.71 nm, respectively. The inner oxide film of milled surface was the thickest while that of electro-polished surface was the thinnest. The standard deviations of the inner layer thickness on the electro-polished, ground and milled surfaces were 1.89, 3.26 and 3.97, indicating that the thickness on the milled surface was the least uniform among three different surfaces. 3.5. Elemental distribution of oxide film The EDX line scan analysis was performed to analyse the elemental distribution of the cross-sectional samples. are the EDX line scan profiles across the oxide films formed on the electro-polished, ground, and milled surfaces, respectively. The vertical coordinate of the EDX result is the count of signal, which represents the quantity of each element. The vertical dashed lines represent the boundaries between each layer, which was established by contrasting the TEM image (Figs. 7–9) and STEM HAADF electron profile (displayed in the top panel of 6

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Fig. 6. SEM images of the surface morphologies of oxide films grown on 316 austenitic stainless steel with different surface states exposed to simulated PWR for 120 h: (a)–(c) electropolished; (d)–(f) ground; (g)–(i) milled.

It is obvious that chromium in the inner oxide film comes from matrix, which would cause the reduction of the quantity of chromium in matrix. However, the relative weight percent profile of chromium beneath the oxide/metal interface was almost unchanged in Fig. 14. As the concentration percent is a relative value, it can’t represent the change of absolute quantity. Nevertheless, allowing for the low diffusion rate of nickel compared to chromium and iron [2], and that chromium and iron have a higher tendency to be dissolved than nickel [27], it is reasonable to consider that the nickel does not diffuse nor reduce at every place in the matrix. In other words, it is reasonable to assume that the quantity of nickel in matrix remains unchanged during oxidation. So we can use the ratio of concentration between chromium and nickel (Cr/Ni) to represent the variation tendency of chromium in matrix beneath oxide film. Fig. 15a shows the variation of the concentration ratio between chromium and nickel along depth. The horizontal dashed line represents the ratio between chromium and nickel concentration in the bulk material. An obvious chromium depletion is seen in Fig. 15a. The quantity of chromium rapidly decreases from matrix to surface. What’s more, the chromium depletion beneath the oxide/metal interface became more and more aggravated from electropolished sample, ground sample to milled sample. Similarly, Fig. 15b shows the variation of the ratio between iron and nickel along depth.

employed as a reference of the depth direction (namely 0 position). The horizontal dashed line represents the relative weight percent of the four metallic elements in the bulk material excluding other elements. In general, the concentration of iron in the outer oxide film was very high, and the inner oxide film was rich in chromium. In addition, molybdenum enrichment was also observed in the inner oxide film. The concentration of nickel reached the highest at the oxide/metal interface and then decreased along the depth direction in the matrix. All these results conform to the EDX line scan analysis. Furthermore, Fig. 14 also shows differences of metallic element concentration among different machined surfaces. In the inner oxide film, the concentration of chromium on the electro-polished surface was higher than that of the ground surface, and the milled sample was the lowest. On the contrary, the concentration of molybdenum on the milled surface was the highest, followed by the ground one, and the electro-polished surface was lowest. At the oxide/metal interface, the concentration of nickel on the milled sample was the highest, followed by the ground and electropolished samples. Meanwhile, in the subsurface the nickel enrichment, which reflects the dealloying of Fe and Cr, was become more and more aggravated with the increase of surface cold working. However, the concentrations of iron in each layer among different finished surfaces were almost similar. 7

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Fig. 7. TEM image of cross-sectional morphology of oxide film grown on electro-polished 316 austenitic stainless steel exposed to simulated PWR. Diffraction pattern showing that the crystal in the outer layer is magnetite (Fe3O4, Orthorhombic) (zone axis [0 −1 1]). Additionally, EDX point analysis region refers to outer layer (A), inner layer (B), interface (C), and matrix (D, E, F with a distance from interface 15, 30, and 70 nm, respectively) are also shown.

Fig. 8. TEM image of cross-sectional morphology oxide film grown on ground 316 austenitic stainless steel exposed to simulated PWR. Diffraction pattern showing that the crystal is magnetite (Fe3O4, Orthorhombic) (zone axis [0 −1 2]). Additionally, EDX point analysis region refers to outer layer (A), inner layer (B), interface (C), and matrix (D, E, F with a distance from interface 15, 40, and 85 nm, respectively) are also shown.

electro-polished sample are given in Fig. 16b–f. OH− and O2− peaks were identified, and the OH− peak is stronger than the O2− peak (Fig. 16b). Chromium was trivalence (Fig. 16c). The peak of iron was mainly Fe3O4 (Fig. 16d). As the contents of Ni and Mo were every low in the oxide film, the signals of both elements were too bad to fit. But the peaks of Ni2+ (Fig. 16e) and Mo6+ (Fig. 16f) were still clear. For the ground and milled surfaces, the detailed spectra of each element in oxide film were provided in the Supplementary material (Figs. S1–S5 of the Supplementary material). Fig. 17 shows the XPS spectra of the oxide films after argon ion sputtering for 300 s. XPS survey spectra of the sputtered oxide film, shown in Fig. 17a, identified O, Cr, Fe, Ni, Mo and C peaks. The detailed spectra of each element in sputtered oxide film are still almost the same for the electro-polished, ground and milled

The more iron depletion was also observed with the increase of cold working, and the most pronounced reduction occurred at the depth less 30 nm of the milled surface.

3.6. Valence states of elements in oxide film XPS analysis was undertaken to provide detailed information on valence states of elements at the unsputtered oxide film and sputtered oxide film. XPS survey spectra of the unsputtered oxide film, shown in Fig. 16a, identified O, Cr, Fe, Ni, Mo and C peaks. It exhibits that the composition of oxide film on three samples is similar. Moreover, high resolution XPS spectra of O1s, Cr2p, Fe2p, Ni2p and Mo3d were conducted. To simplify the illustration, only the XPS detailed spectra of the 8

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Fig. 9. TEM image of cross-sectional morphology oxide film formed on milled 316 austenitic stainless steel exposed to simulated PWR. Diffraction pattern showing that the crystal is magnetite (Fe3O4, Orthorhombic) (zone axis [0 −1 1]). Additionally, EDX point analysis region refers to: outer layer (A), inner layer (B), interface (C), matrix (D, E, F with a distance from interface 20, 60, and 80 nm, respectively) are also shown.

4. Discussion

surfaces (Figs. S6–S10 of Supplementary material). The XPS detailed spectra of the sputtered electro-polished sample are given in Fig. 17b–f. Except for the valence states of elements detected by the previous analysis before sputtering, the lower valence states were observed. In Fig. 17e only Ni0 was observed. Fe2+ increased a lot, and a small amount of Fe0 was found (Fig. 17d). For molybdenum, Mo6+, Mo4+ and Mo0 were found. No matter the unsputtered oxide film or the sputtered oxide film, the XPS results show that the surface machining operations do not affect the valence state of elements in the oxide films.

4.1. Oxide film characterization The morphologies of oxide films on three machined surfaces are similar. The oxide film consists of an outer layer and an inner layer. The outer layer is composed of iron-rich large oxide particles, and TEM observation suggested that the oxide particle consists of magnetite (Fe3O4). The inner layer is composed of compact nanocrystalline which is rich in chromium. This kind of morphology of oxide film is in agreement with other studies [1–8]. The distribution of iron and chromium in oxide film could be related to their different affinities for oxygen. At the beginning of the 316 austenite stainless steel exposed to PWR primary water environment, Fig. 10. Frequency distribution histograms of the inner oxide film thickness grown on 316 austenitic stainless steel with different surface states: (a) electro-polished; (b) ground; (c) milled.

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Fig. 11. EDX line scan profiles across the oxide film grown on electropolished surface exposed to simulated PWR for 120 h.

the protective oxide film has not yet formed on the surface, and all the elements in the steel have equal opportunity to contact with oxygen and to produce oxide. As the oxygen affinity for chromium is larger than those of iron and other elements in alloy. The oxygen is more ready to react with chromium than iron and nickel, which leads to the formation of a chromium-rich inner oxide layer. The stability of chromium oxides determines that it will be a compact and protective film. When the thickness of this film reaches a certain value, it will prevent metal from directly contacting with oxygen. Then the oxidation rate will become slow. As a result, the contact of the metal cations with oxygen can only be realized by means of the diffusion of metal cations from matrix to solution, and the diffusion of oxygen from solution to matrix through oxide film. As the diffusion rate of iron is faster than those of chromium and nickel [1], the high concentration of iron at the oxide/metal interface will drive iron ion to diffuse from matrix to solution through oxide film. Then iron ion reacts with oxygen at the solution/oxide interface and forms the iron-rich outer oxide film with a high valence state of Fe3O4. Although most of the studies reported that the outer oxide layer was rich in iron, there is a divergence in what the outer oxide is. Some studies tended to magnetite (Fe3O4) [2,6] while others thought it was iron–nickel spinel oxide [1,3], or even both of them [8]. In the present study only magnetite (Fe3O4) was observed in the outer layer by electron diffraction. In addition, there is also a divergence in the morphology of outer layer. de Jesús et al. [32] reported that the outer oxide layer of 316 stainless steel was composed of large particles (nickel-enriched Fe3O4-type structure) and intermediate or small particles (α-Fe2O3 but chromium-enriched Fe3O4 type in only hydrogen water chemistry) at 288 °C water in cyclic normal and hydrogen water

chemistries for several weeks. Cissé et al. [8] also found the large particles (Fe3O4 magnetite type) and small particles (Fe(Cr/Ni)2O4 spinel type) in outer layer of 304L austenitic stainless steel in simulated PWR primary water at 350 °C for 500 h. However, in the present study the particularly large particle was not observed. The test environment in this study is similar to Cissé et al.’s, so the test environment can’t be the main reason for the formation of the particularly large particle. The size of the particularly large particle in Cissé’s study was over 1 micrometer, which is much bigger than the oxide particle with a size of only several hundred nanometres in our study. In addition to material difference, the immersion time of 120 h in the present study is far short than Cissé’s study. It is reasonable to conclude that the immersion time in this study was too short to allow the oxide particle grow into micron size. Nickel enrichment at the oxide/metal interface was observed in our study. This result is consistent with the studies reported previously [2,4,6]. Due to the lowest oxygen affinity of nickel among all the elements in 316 austenite stainless steel [27], nickel is the least likely to be oxidized. In addition, the diffusion rate of nickel is slow compared with iron and chromium [2]. As a result, chromium with the highest oxygen affinity is oxidized to form the inner oxide film and iron diffuse from matrix to solution through inner oxide film to form the outer oxide film, and then nickel will be left at the interface between oxide and matrix. In addition, our study shows that the concentration of molybdenum in the inner oxide film was relatively high. To the best of my knowledge, molybdenum enrichment in the inner oxide film has rarely been mentioned in studies of 316 austenitic stainless steel in simulated PWR water. Terachi et al. [2] who studied oxide film for 316 stainless steel in 10

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Fig. 12. EDX line scan profiles across the oxide film grown on ground 316 austenitic stainless steel exposed to simulated PWR for 120 h.

may result in the reduction of metal oxide [29–31]. Argon ion sputtering induces many changes in composition, and in oxidation state. So there is a limitation for the present study of analysis with argon ion sputtering. Combining with above analysis, there are three possible reasons for the existence of lower valence state: (1) exist originally in the oxide film, (2) released from the specimen, and (3) the high valence state elements are reduced by the argon ion sputtering. Despite the limitation of analysis with argon ion sputtering, it still can be concluded that surface machining operations do not affect the valence state of elements in the oxide film from the present study.

simulated PWR primary water at 320 °C particularly mentioned that molybdenum was not found in the oxide film. However, in the present study, both in EDX analysis and XPS, molybdenum enrichment in the inner oxide film was observed in all samples despite the fact that the concentration of molybdenum in the polished surface was lower than those of the machined surfaced. The present study shows that the lower valence state of element was detected after sputtering for 300 s. One of the possible reasons why lower valence state elements were detected after sputtering is that the oxygen potential decreased with the depth. The low oxygen potential at the oxide/metal interface resulted in the instability of iron oxides [28], let alone for Ni and Mo whose oxygen affinities are lower than iron. Therefore, it is possible that the lower valence state elements exist originally in the oxide film. XPS depth profile with argon ion sputtering was previously used in many oxide film studies, and the lower valence states of elements were commonly observed after sputtering [7,18–20]. The majority of the researchers did not explain whether the lower valence state elements existed originally or not. Han et al. [19] suggested that X-ray could penetrate the surface film, which resulted in the detection of metallic state. As the sputtering rate for corrosion oxide spinel is lower than that for Ta2O5 [23], for the present study the calculated removed thickness by sputtering was less than 3 nm. The analysis depth of the XPS system is 7 nm, and the inner oxide film is more than 20 nm. As a result, in this study the signal of lower valence state element came from the oxide film rather than the matrix. Guo et al. [18] thought that metallic Ni was released from the specimen. However, numerous researches also pointed out that argon ion sputtering

4.2. Effects of surface machining on oxide film characterization Compared three different surfaces, the surface Vickers micro-hardness increased in turn from the electro-polished surface, ground surface to milled surface. For austenitic stainless steels, the change of surface micro-hardness due to surface machining is mainly owing to two reasons [33]: (1) the increase in dislocation density, and (2) the phase transformation of austenite to martensite. In this study, the phase transformation was not observed from X-ray diffraction profiles. So the increase in surface micro-hardness was mainly owing to the change of dislocation density. There is a positive correlation between microhardness and dislocation density [34], it implied that the dislocation density increased from the electro-polished surface, ground surface to milled surface. The grain size measured by XRD seems not to conform to the result which directly measured on the TEM image. This may be attributed to 11

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Fig. 13. EDX line scan profiles across the oxide film grown on milled 316 austenitic stainless steel exposed to simulated PWR for 120 h.

thicker inner film and a higher concentration of nickel at interface. Similar to our observation, Berge [36] reported that the oxidation kinetics of 304 austenitic stainless steel with electro-polishing was three to four times slower than that with mechanical finishing treatment in water at 300 °C. Ziemniak et al. [3,23] also reported that the corrosion rate of machined 304 austenitic stainless steel was far above electropolished one in mildly alkaline, hydrogenated water at 260 °C. However, Cissé et al. [8] studied the effect of surface preparation on the corrosion of austenitic stainless steel 304L in high temperature steam and simulated PWR primary water with no grain size in the subsurface provided. They reported that the Cr-rich inner layer on polished sample was thicker than ground one. This seems contrary to our study. It may be due to different ways of surface polishing and grinding. Their ground sample was prepared by precision grinding, and polished sample was ended with colloidal silica suspension polishing. In the present study, the electrochemical polishing was applied to remove the Beilby layer, and the grinding was carried out by a conventional grinding machine. Moreover, the uniformity of the inner layer thickness became poor from electro-polished surface, ground surface to milled surface. One of the reasons is possibly due to the different surface roughness among three different surfaces. Li et al. [16,17] thought that the “peak” of the fluctuating surface had a higher tendency to corrode than the “valley” on rough surface. The different corrosion rate between local regions may result in the different film thickness between local regions. In other words, there are more metal cations diffusing to the surface and reacting with oxygen at “peaks” than that at “valleys”. Since the surface

the different measuring depths. The cold worked layer was not homogeneous but a layer in which microstructure varies along depth. The Xray could pass through the material for several micrometres, while the TEM only observed about 20 nm below the oxide film. Despite those differences, it still can be inferred that the grain size of the electropolished surface was the largest, and that of the milled surface was the smallest. Namely, the density of grain boundaries, whose trend was consistent with that of dislocation density, increased gradually from the electro-polished surface, ground surface to milled surface. After PWR primary water exposure, the milled surface formed the thickest oxide film, the highest concentration of molybdenum in the inner layer, and the highest concentration of nickel at the oxide/metal interface. On the contrary, the electro-polished sample resulted in the thinnest oxide film with the least molybdenum, and the concentration of nickel at the interface of electro-polished surface was the lowest. These probably were attributed to the differences of dislocation density and grain size among the different surfaces. The diffusion rate of ion in grain boundary is largely faster than that in crystal lattice. The dislocations and grain boundaries can act as the short circuits for the outward diffusion of metal cations and the inward diffusion of oxygen [35]. So in the surface cold worked sample, more metal cations were applicable to the formation of the inner oxide film at the initial stage of oxidation, which led to a thicker film with more molybdenum. At the same time, more chromium and iron diffused out of matrix, which made the concentration of nickel at the oxide/metal interface increased more greatly. As a result, the more heavily cold working leads to a

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Fig. 14. Relative weight percent profiles across the oxide film obtained by EDX point analysis. Horizontal dashed line represents the relative weight percent of the four metallic elements in the bulk material excluding other elements.

Fig. 15. Cr/Ni concentration ratio and Fe/Ni concentration ratio varying along depth. (a) Cr/Ni concentration ratio (b) Fe/Ni concentration ratio.

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Fig. 16. XPS spectra of the unsputtered oxide films. (a) XPS survey spectra for all elements on three samples without sputtering. (b) Detailed spectra of O1s for unsputtered electropolished sample. (c) Detailed spectra of Cr2p for unsputtered electro-polished sample. (d) Detailed spectra of Fe2p for unsputtered electro-polished sample. (e) Detailed spectra of Ni2p for unsputtered electro-polished sample. (f) Detailed spectra of Mo3d for unsputtered electro-polished sample.

A lower concentration of chromium in the inner oxide film of the cold worked surfaces were observed compared to the polished surface in this study. Although the increased dislocations and grain boundaries in the cold worked layer can promote outward diffusion of metal cation to be oxidized, it seems that the promotion is greater for Fe and Ni than Cr. Fig. 18 shows the Fe/Cr ratio and Ni/Cr ratio in the inner oxide film for three surfaces. The Fe/Cr ratio increased from the electro-polished surface, ground surface to milled surface, which indicated that the machining -induced promotion is greater for Fe than Cr. Similarly, the promotion is greater for Ni than Cr. Therefore the more heavily cold working led to a lower concentration of chromium in the inner oxide film.

roughness of milled sample was the highest, which would lead to a larger difference of corrosion rate between local regions. As a result, the uniformity of the inner layer thickness was the poorest. The results in Fig. 14 showed that the concentrations of iron in each layer of the oxide film were almost similar among three different surfaces. Despite the fact that the dislocations and grain boundaries can promote outward diffusion of iron cation into inner oxide film, the oxygen is more ready to react with chromium than iron in the inner oxide film where the oxygen potential is low. The redundant iron can only diffuse outwards the solution/oxide interface and form the outer oxide film. As a result, the concentration of iron in inner oxide film was not obviously affected by the surface cold working. Instead, the outermost oxide particle (outer oxide film) is more and larger for the cold worked sample. As the outer oxide film was observed composed of magnetite (Fe3O4), the concentrations of iron in outer oxide film were similar among three surfaces. 14

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Fig. 17. XPS spectra of the oxide films after sputtering for 300 s. (a) XPS survey spectra for all elements on three samples after sputtering. (b) Detailed spectra of O1s for sputtered electropolished sample. (c) Detailed spectra of Cr2p for sputtered electro-polished sample. (d) Detailed spectra of Fe2p for sputtered electro-polished sample. (e) Detailed spectra of Ni2p for sputtered electro-polished sample. (f) Detailed spectra of Mo3d for sputtered electro-polished sample.

4.3. Chromium depletion beneath the oxide/metal interface due to surface cold working In addition to change of relative concentration of chromium in the inner oxide film, the cold worked surfaces also caused more aggravated chromium depletion in terms of Cr/Ni ratio beneath the oxide/metal interface. During the oxidation, chromium in the inner oxide film are being dissolved in the solution all the time and the inner layer is also growing. For purpose of sustaining the inner oxide film, the chromium beneath the oxide/metal interface diffuses into the oxide film to replenish chromium continuously. The movement of chromium in matrix is realized by diffusion. As the corrosion kinetics of austenitic stainless steel was observed to be parabolic by Ziemniak and Hanson [23]. The theoretical description of chromium profiles in subsurface can be given by the equation [37]: Fig. 18. Fe/Cr concentration ratio and Ni/Cr concentration ratio in the inner oxide film for three surfaces.

C (z , t ) = C0 + (CA − C0)erf( 15

z ) 2 D ⋅t

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x

2

Where erf(x) is Gauss error function.erf(x ) = π ∫0 e−η dη , C(z, t) the chromium concentration at a distance z from the oxide/metal interface at time t, C0 the concentration of chromium on the surface of matrix at the oxide/metal interface, CA the bulk chromium concentration, D the diffusion coefficient of chromium. It is obvious that Gauss error function is an increasing function. When the time t and the depth z is constant, the increase of diffusion coefficient will result in the decrease of concentration of chromium. It was mentioned above that the dislocation density and grain boundary density of surface cold worked sample was larger than that of electropolished one. The diffusion of species through grain boundaries and dislocation is faster than that through crystal lattice. So the chromium diffusion coefficient of the cold worked surface was larger than that of the electro-polished surface, which led to the unduly chromium depletion beneath the oxide/metal interface in cold worked sample.

5. Conclusions 1. The morphologies and compositions of oxide films on the three kinds of surfaces are almost similar. They are comprised of a twolayer structure. The outer layer is single oxide particle of magnetite (Fe3O4), and the inner layer is composed of chromium-rich nanocrystalline. The valence state of elements in the oxide film was not influenced by the surface machining operations. 2. The uniformity of the inner layer thickness was affected by surface preparation methods. The oxide film on the electro-polished surface exhibited the best uniformity, followed by the ground surface and the milled surface. The mean thickness of inner layer increased from the electro-polished surface, ground surface to milled surface. 3. The concentration of chromium in the inner oxide film was the highest on the electro-polished surface followed by the ground, and the lowest from the milled surface. Instead, the highest concentration of molybdenum in the inner layer was observed on the milled surface, and the lowest one was from the electro-polished surface. The concentrations of iron in each layer of the oxide films were almost similar among different surfaces. 4. Nickel enrichment was observed at the oxide/metal interface for all surfaces. The concentration of nickel at the oxide/metal interface was the highest in the milled surface, followed by ground one and electro-polished one. 5. The more heavily cold worked layer results in more aggravated chromium depletion and iron depletion beneath oxide film.

4.4. Implication for SCC According to the point defect model [38,39], the oxide film grows by migration of metal cation at matrix from oxide/metal interface to metal/solution interface. At the same time, vacancies would be left at oxide/metal interface and finally lead to a local accumulation, which results in stresses in the passive film and its subsequent breakdown. In the present study, the nickel enrichment at the oxide/metal interface was observed. Moreover, the surface cold worked layer led to a serious nickel enrichment. It is speculated that the surface machining would lead to more accumulation of vacancies in the subsurface. Yang et al. [40] found that thick film increased susceptibility to stress corrosion cracking of type 304 stainless steel in high temperature waters. The present results showed that cold worked surface led to not only a thick oxide film but also a poor uniformity of thickness. From the viewpoint of mechanics, the nonuniform thickness of film may result in stress concentration in the presence of external loading or machining-induced residual stress. So the thick and nonuniform oxide film resulted from machining has a low fracture resistance. Actually, our previous study [41] showed that the residual tensile stress remained in the machiningaffected layer has a great influence on the initiation of surface microcracks. Breummer and Thomas [42] found that the nature crack generated in service in BWR during long time was very different from the crack generated in the laboratory by crack growth rate tests. The nature crack tips showed the presence of locally “dealloyed” zones of Fe/Cr-depleted, Ni-rich metal. Especially, the Cr reduced by 78.9 percent compared with the bulk material. Staehle [43] reviewed the TEM studies of stress corrosion cracking in Fe-Cr-Ni alloys used in water cooled nuclear power plants in the range of 250–350 °C with hydrogenated and oxygenated water and with some contaminated environments, he found that there was a very clear dealloying occurring in the metal ahead of the crack especially in the stainless steels. These were similar to the oxidation behaviour of stainless steel with surface cold working discussed in this study. Our study shows that the surface cold worked layer leads to low chromium concentration in the inner oxide layer and unduly chromium depletion beneath the oxide layer. Numerous studies have confirmed chromium is the key element to form a compact and protective oxide film. This roughly suggests that oxide film grown on the machined 316 stainless steel is less protective than the film formed on polished sample. The chromium depletion at the oxide/metal interface may lead to a weak adhesion of the oxide film to bulk material, causing an exfoliation of the passive protective film. Furthermore, if the surface oxide film cracks, the matrix beneath the oxide film will be exposed to etchant solution. In addition to machining-induced residual stress, the chromium depletion beneath the oxide/metal interface due to surface cold working would be the impetus of crack propagation into bulk material.

Acknowledgments This work is supported by the National Natural Science Foundation of China (Grant No. 51375182). The authors thank Analytical and Testing Center of HUST for FSEM and FTEM measurements, and the State Key Laboratory of Materials Processing and Die & Mould Technology of HUST for XRD measurement. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.corsci.2017.06.019. References [1] B. Stellwag, The mechanism of oxide film formation on austenitic stainless steels in high temperature water, Corros. Sci. 40 (1998) 337–370. [2] T. Terachi, K. Fujii, K. Arioka, Microstructural characterization of SCC crack tip and oxide film for SUS 316 stainless steel in simulated PWR primary water at 320 C, J. Nucl. Sci. Technol. 42 (2005) 225–232. [3] S.E. Ziemniak, M. Hanson, P.C. Sander, Electropolishing effects on corrosion behavior of 304 stainless steel in high temperature hydrogenated water, Corros. Sci. 50 (2008) 2465–2477. [4] S. Ghosh, M.K. Kumar, V. Kain, High temperature oxidation behavior of AISI 304L stainless steel–effect of surface working operations, Appl. Surf. Sci. 264 (2013) 312–319. [5] M.C. Sun, X.Q. Wu, Z.E. Zhang, E.H. Han, Oxidation of 316 stainless steel in supercritical water, Corros. Sci. 51 (2009) 1069–1072. [6] R. Soulas, M. Cheynet, E. Rauch, T. Neisius, L. Legras, C. Domain, Y. Brechet, TEM investigations of the oxide layers formed on a 316L alloy in simulated PWR environment, J. Mater. Sci. 48 (2013) 2861–2871. [7] W.J. Kuang, X.Q. Wu, E.H. Han, The oxidation behaviour of 304 stainless steel in oxygenated high temperature water, Corros. Sci. 52 (2010) 4081–4087. [8] S. Cissé, L. Laffont, B. Tanguy, M. Lafont, E. Andrieu, Effect of surface preparation on the corrosion of austenitic stainless steel 304L in high temperature steam and simulated PWR primary water, Corros. Sci. 56 (2012) 209–216. [9] K.D. Ralston, N. Birbilis, Effect of grain size on corrosion: a review, Corrosion 66 (2010) 75005. [10] A. Di Schino, J.M. Kenny, Effects of the grain size on the corrosion behavior of refined AISI 304 austenitic stainless steels, J. Mater. Sci. Lett. 21 (2002) 1631–1634. [11] J. Lv, H. Luo, Comparison of corrosion behavior between coarse grained and nano/ ultrafine grained 304 stainless steel by EWF, XPS and EIS, J. Nucl. Mater. 452 (2014) 469–473. [12] T.S. Li, L. Liu, B. Zhang, Y. Li, F.K. Yan, N.R. Tao, F.H. Wang, Passive behavior of a bulk nanostructured 316L austenitic stainless steel consisting of nanometer-sized grains with embedded nano-twin bundles, Corros. Sci. 85 (2014) 331–342.

16

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S. Wang et al. [13] G.T. Burstein, P.C. Pistorius, Surface roughness and the metastable pitting of stainless steel in chloride solutions, Corrosion 51 (1995) 380–385. [14] G.T. Burstein, S.P. Vines, Repetitive nucleation of corrosion pits on stainless steel and the effects of surface roughness, J. Electrochem. Soc. 148 (2001) B504–B516. [15] M.B. Leban, Mikyška, T. Kosec, B. Markoli, J. Kovač, The effect of surface roughness on the corrosion properties of type AISI 304 stainless steel in diluted NaCl and urban rain solution, J. Mater. Eng. Perform. 23 (2014) 1695–1702. [16] W. Li, D.Y. Li, Variations of work function and corrosion behaviors of deformed copper surfaces, Appl. Surf. Sci. 240 (2005) 388–395. [17] W. Li, D.Y. Li, Influence of surface morphology on corrosion and electronic behavior, Acta Mater. 54 (2006) 445–452. [18] Y.L. Guo, E.H. Han, J.Q. Wang, Effects of surface states on the oxidation behavior of 316LN stainless steel in high temperature pressurized water, Mater. Corros. 66 (2015) 670–680. [19] Y. Han, J. Mei, Q. Peng, E. Han, W. Ke, Effect of electropolishing on corrosion of nuclear grade 316L stainless steel in deaerated high temperature water, Corros. Sci. 112 (2016) 625–634. [20] H. Ming, Z. Zhang, J. Wang, R. Zhu, J. Ding, J. Wang, E. Han, W. Ke, Effect of surface state on the oxidation behavior of welded 308L in simulated nominal primary water of PWR, Appl. Surf. Sci. 337 (2015) 81–89. [21] G. Han, Z. Lu, X. Ru, J. Chen, Q. Xiao, Y. Tian, Improving the oxidation resistance of 316L stainless steel in simulated pressurized water reactor primary water by electropolishing treatment, J. Nucl. Mater. 467 (2015) 194–204. [22] M. Warzee, J. Hennaut, M. Maurice, C. Sonnen, J. Waty, P. Berge, Effect of surface treatment on the corrosion of stainless steels in high-temperature water and steam, J. Electrochem. Soc. 112 (1965) 670–674. [23] S.E. Ziemniak, M. Hanson, Corrosion behavior of 304 stainless steel in high temperature hydrogenated water, Corros. Sci. 44 (2002) 2209–2230. [24] G.K. Williamson, W.H. Hall, X-ray line broadening from filed aluminium and wolfram, Acta Metall. 1 (1953) 22–31. [25] M. Imran, P.T. Mativenga, A. Gholinia, P.J. Withers, Evaluation of surface integrity in micro drilling process for nickel-based superalloy, Int. J. Adv. Manuf. Technol. 55 (2011) 465–476. [26] M. Lundberg, J. Saarimäki, J.J. Moverare, M. Calmunger, Surface integrity and fatigue behaviour of electric discharged machined and milled austenitic stainless steel, Mater. Charact. 124 (2017) 215–222. [27] X. Ren, K. Sridharan, T.R. Allen, Corrosion of ferritic-martensitic steel HT9 in supercritical water, J. Nucl. Mater. 358 (2006) 227–234. [28] Y. Kim, Characterization of the oxide film formed on type 316 stainless steel in

[29]

[30]

[31] [32] [33]

[34] [35]

[36] [37] [38] [39]

[40]

[41]

[42]

[43]

17

288 C water in cyclic normal and hydrogen water chemistries, Corrosion 51 (1995) 849–860. M. Ueda, M. Nanko, K. Kawamura, T. Maruyama, Formation and disappearance of an internal oxidation zone in the initial stage of the steam oxidation of Fe–9Cr–0.26 Si ferritic steel, Mater. High Temp. 20 (2003) 109–114. K.S. Kim, W.E. Baitinger, J.W. Amy, N. Winograd, ESCA studies of metal-oxygen surfaces using argon and oxygen ion-bombardment, J. Electron Spectrosc. 5 (1974) 351–367. R. Holm, S. Storp, E SCA studies on changes in surface composition under ion bombardment, Appl. Phys. A Mater. Sci. Process. 12 (1977) 101–112. J.C. de Jesús, P. Pereira, J. Carrazza, F. Zaera, Influence of argon ion bombardment on the oxidation of nickel surfaces, Surf. Sci. 369 (1996) 217–230. H. Ding, Y.C. Shin, Multi-physics modeling and simulations of surface microstructure alteration in hard turning, J. Mater. Process. Technol. 213 (2013) 877–886. M.S. De Guzman, G. Neubauer, P. Flinn, W.D. Nix, The Role of Indentation Depth on the Measured Hardness of Materials, Cambridge Univ. Press, 1993 pp. 613. Z. Zhang, J. Wang, E. Han, W. Ke, Characterization of different surface states and its effects on the oxidation behaviours of Alloy 690TT, J. Mater. Sci. Technol. 28 (2012) 353–361. P. Berge, Importance of surface preparation for corrosion control in nuclear power stations, Mater. Perform. 36 (1997). H.S. Carslaw, J.C. Jaeger, Conduction of Heat in Solids, 2nd ed., Clarendon Press, Oxford, 1959. C.Y. Chao, L.F. Lin, D.D. Macdonald, A point defect model for anodic passive films I. Film growth kinetics, J. Electrochem. Soc. 128 (1981) 1187–1194. L.F. Lin, C.Y. Chao, D.D. Macdonald, A point defect model for anodic passive films II. Chemical breakdown and pit initiation, J. Electrochem. Soc. 128 (1981) 1194–1198. W. Yang, G. Zhao, M. Zhang, J. Congleton, An AES investigation of the surface films formed on stress corrosion test specimens of type 304 stainless steel in high temperature water, Corros. Sci. 33 (1992) 89–102. W. Zhang, K. Fang, Y. Hu, S. Wang, X. Wang, Effect of machining-induced surface residual stress on initiation of stress corrosion cracking in 316 austenitic stainless steel, Corros. Sci. 108 (2016) 173–184. S.M. Bruemmer, L.E. Thomas, Mechanistic Insights on Environmental Degradation From Nanometer-Scale Crack-Tip Measurements, Pacific Northwest National Laboratory (PNNL), Richland, WA (US), 2005. R.W. Staehle, Critical analysis of tight cracks, Corros. Rev. 28 (2010) 1–103.