Effect of surface treatments on microstructure and stress corrosion cracking behavior of 308L weld metal in a primary pressurized water reactor environment

Journal Pre-proof Effect of surface treatments on microstructure and stress corrosion cracking behavior of 308L weld metal in a primary pressurized water reactor environment Lijin Dong (Investigation) (Validation) (Data curation) (Writing original draft) (Writing - review and editing), Xiaolong Zhang (Investigation) (Visualization), Yaolei Han (Investigation), Qunjia Peng (Supervision) (Writing - review and editing), Ping Deng (Investigation), Shuliang Wang (Writing - review and editing)

PII:

S0010-938X(19)32236-X

DOI:

https://doi.org/10.1016/j.corsci.2020.108465

Reference:

CS 108465

To appear in:

Corrosion Science

Received Date:

24 October 2019

Revised Date:

8 January 2020

Accepted Date:

9 January 2020

Please cite this article as: Dong L, Zhang X, Han Y, Peng Q, Deng P, Wang S, Effect of surface treatments on microstructure and stress corrosion cracking behavior of 308L weld metal in a primary pressurized water reactor environment, Corrosion Science (2020), doi: https://doi.org/10.1016/j.corsci.2020.108465

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Effect of surface treatments on microstructure and stress corrosion cracking behavior of 308L weld metal in a primary pressurized water reactor environment Lijin Donga,*, Xiaolong Zhanga, Yaolei Hanb, Qunjia Pengb,*, Ping Dengc, Shuliang Wanga a

School of School of Materials Science and Engineering, Southwest Petroleum University, Chengdu 610500, China Suzhou Nuclear Power Research Institute, Suzhou 215004, China

c

Nuclear Power Institute of China, Chengdu 610041, China

* Corresponding author. [email protected] (L.J. Dong)

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E-mail address:

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b

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[email protected] (Q.J. Peng)

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Highlights  Milling, mechanically grinding and electropolishing promote SCC of 308L weld metal  Composition change and etched phase boundary cause SCC of electropolished surface  Machining produces a nanocrystalline layer and an underlying deformed layer  Elongated ferrite phases in nanocrystalline layer increase SCC of machined surface

ABSTRACT

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Effect of surface treatments on microstructure and stress corrosion cracking (SCC)

of 308L weld metal in primary pressurized water reactor environment was evaluated by microstructure characterization and bent beam tests. The results showed that no SCC was observed on colloidal silica slurry polished surface while the change in composition and etched phase boundary promote SCC initiation after electropolishing. Further, milling and grinding produce a nanocrystalline layer and an underlying deformed layer on the surface. The nano-sized δ ferrite phase has a detrimental effect 1

on SCC due to the synergistic effect of high concentration of dislocations and its nano size.

Keywords: A. stainless steel; B. SEM; B. TEM; C. Stress corrosion; C. Welding

1. Introduction In the primary circuit of pressurized water reactor (PWR) nuclear power plants, stainless steels and their weld meals have been widely used as key structural material

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owing to the high resistance to corrosion and satisfactory mechanical properties. However, the stainless steels were found to be susceptible to stress corrosion cracking

(SCC) in primary PWR environments [1-12]. Since milling, local grinding, shot peening and polishing alter the microstructure, topography and composition of the

treatments can play an important role in SCC.

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material on surface that is in direct contact with high temperature water, surface

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Machining such as milling and grinding is widely-used to produce a desired geometry, but could generate residual stresses and change the microstructure of the

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material. To understand the effect of surface machining on microstructure and SCC behavior of stainless steels, a number of studies have been conducted in laboratory [1,

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10, 13-20], and revealed that the surface layer of solution annealed stainless steel after machining could be divided into a nanocrystalline layer and an underlying deformed layer. When exposed in primary PWR environments, the occurrence of fast oxidation

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in the nanocrystalline layer resulted in the formation of a high-porosity, less protective

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oxide film [3, 13]. Cracks could easily initiate in the oxidized nanocrystalline layer, although only a limited number of them penetrated into the underlying deformed layer [3, 13, 14]. Other investigations on warm-forged 304L stainless steel [15, 16], however, revealed the nanocrystalline layer introduced by milling and grinding could increase the SCC initiation resistance because the high density of grain boundaries and the highly deformed microstructure promoted the formation of a more uniform oxide film in the nanocrystalline layer [16]. 2

Electropolishing changes the residual strain, roughness and composition on the surface of austenitic alloys, and therefore can affect the SCC in high temperature water [20-27]. A couple of studies concluded that the electropolishing increased the SCC susceptibility because it decreased the compressive stress at the surface compared to mechanical polishing [21, 22]. However, other studies found that electropolishing could retard SCC initiation since it lowered the surface residual strain [23-26]. Further, a more recent investigation revealed that the electropolishing induced Cr enrichment at the grain boundary, promoted the formation of a protective oxide in primary PWR

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environments, and therefore decreased the SCC susceptibility of Alloy 182 [27]. While many studies have investigated the effect of surface treatments on the SCC behavior of austenitic alloys in primary PWR environments [5, 6, 13-17], only a few

focused on the stainless steel weld metals such as type 308L [18, 28] , which has 3 vol.%

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to 12 vol.% δ ferrite phase to enhance hot cracking resistance [29-31]. The 308L stainless steel weld metal is used to clad the inner surface of reactor pressure vessel and

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attach the safe end of primary circuit piping to reactor pressure vessel nozzle. Due to the Cr-enriched δ ferrite phase along the dendritic grain boundaries, the weld metals

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generally exhibited low stress corrosion crack growth rate in primary PWR environments [7, 32-34]. However, it was reported that cracks initiated from the

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cladding surface of reactor pressure vessels [35, 36]. Further, several studies of the effect of surface treatments on corrosion behavior in high temperature water indicated that both the mechanical polishing and electropolishing changed the thickness and

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protectiveness of the oxide film formed on 308L weld metal [37, 38]. This suggests that surface treatments could alter the SCC behavior of the 308L weld metal, since the

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protectiveness of the surface oxide film is a key factor in SCC. In fact, it was found that SCC initiation of weld metals was promoted by surface grinding in high temperature water [18]. Further, due to differences in microstructure, the effect of surface treatments on SCC of 308L weld metal may differ from 304L and 316L stainless steels, which have been extensively studied [5, 10, 13-16, 39]. For example, the difference in chemical composition and mechanical properties for the δ ferrite and austenite phases 3

in the bulk material could lead to a more complex microstructure in the surface layer after surface treatments, and thus different SCC. In addition, the role of the δ ferrite phase in SCC of weld metals after surface treatments is not well understood. Therefore, investigating the effect of surface treatments on the SCC behavior of 308L stainless steel weld metal in primary PWR environments is required, to clarify the SCC mechanism and develop technologies for mitigating SCC using surface treatments. To clarify different methods of surface finishing on SCC of 308L weld metal, four kind of surface treatments were evaluated in the current study: 40 nm colloidal silica

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slurry polishing, mechanically grinding with dry abrasive paper, milling, and electropolishing. Surface grinding or milling are often the final process for the

fabrication of weld joints, leading to a deformation layer on surface. Electropolishing

and slurry polishing were used to remove the deformation layer although these methods

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were not applied in actual PWRs. Compared to the electropolishing, slurry polishing

removes the surface residual strain but leads to few composition changes on surface.

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As such, the microstructure characterization and evaluation of the SCC behavior after four kind of surface treatments could clarify the effect of microstructure, residual stress

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and surface composition. The microstructure of the 308L weld metal after surface treatments was analyzed by scanning electron microscope (SEM), electron backscatter

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diffraction (EBSD), Vickers microhardness, X-ray photoelectron spectroscopy (XPS), three-dimensional surface morphology analyzer, nanoindenter and transmission electron microscope (TEM). The SCC behavior were evaluated using bent beam tests

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in a simulated primary PWR environment. This method is particularly suited to

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investigating the effect of surface treatments on SCC [1, 27]. 2. Experimental 2.1 Material and specimen As schematically shown in Fig. 1, the 308L stainless steel was machined from the cladding of a mockup of safe-end weld joint to prepare specimens. The mockup is an assembly of a SA508 pipe welded to a 316L safe-end by 308L and 309L filler metals. After the SA508 pipe with an internal diameter of about 600 mm was preheated at 4

temperatures higher than 150 C, a buttering layer of 309L and 308L cladding layer was sequential deposited on the steel (120–180 A, 10–12 V, interpass temperatures for 309L and 308L are less than 250 C and 100 C, respectively). Then post-weld heat treatment at a temperature below 400 °C for about 4 h was performed. The cladded pipe and 316L safe-end were connected by a multi-layer and multi-pass shielded metal arc welding, followed by heating at 600–630 C for about 2 h to relieve the residual stress. The simulated post-weld heat treatment of the weld joint is heated at 600-620 ℃ for about 40 h. The thickness of the 309L buttering and 308L cladding was about 2 mm and 6

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mm, respectively. The chemical compositions of the 308L welding wire and 308L cladding were determined by chemical analysis method, as listed in Table 1.

To clarify the effect of surface treatments on microstructure and SCC behavior of 308L weld metal, the four surface treatments, i.e., colloidal silica slurry polishing,

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mechanically grinding, milling and electropolishing were used to prepare the specimens (Fig. 1). All of the specimens were exacted in the same orientation, i.e., their

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longitudinal direction was parallel to the axial direction of the safe-end. The mechanically ground specimens were prepared using 80 grit dry abrasive paper of

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diameter 300 mm pasted on a disc. The rotating speed of the disc was 500 r/min. After mechanically grinding, the cutting depth of was about 0.02 mm. Some of the specimens

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were further ground with 240, 400, 800, 1200 and 2000 grit waterproof abrasive paper and followed by polishing using 1 μm diamond paste. They were then electropolished or manually polished by 40 nm colloidal silica slurry. The electropolishing was

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conducted in an electrolyte containing 80 vol.% perchloric acid (70 wt.%) and 20 vol.% glacial acetic acid (98 wt.%) at −20 °C, 0.2 A/cm2 for 20 s. The removed material

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thickness was about 10 μm. The 40 nm colloidal silica slurry polishing employed a period of >3 h to ensure the removal of surface residual strain. The milled specimens were prepared by machining the surface of the cladding using a milling cutter. The cutting speed and depth were 150 m/min and 1 mm, respectively, and the feed rate was 0.3 mm/rev. After different surface treatments, all these specimens used for the SCC test were reduced to a dimension of 40 mm × 10 mm × 2 mm by wire electrical 5

discharge machining and mechanically grinding on the opposite side of the surface. The longitudinal direction of the milled specimens for SCC tests was perpendicular to the milling marks. The δ ferrite phase in 308L cladding was island-shaped, as observed on the surface of the cladding by optical microscope (OM), shown in Fig. 2a. However, on the crosssection of the cladding, vermicular δ-ferrite with a width of less than 5 μm was observed at the dendritic grain boundaries of the austenite matrix (Fig. 2b). The growth direction of the dendritic grain was approximately perpendicular to the surface. Further, carbides

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with a semi-continuous distribution along the γ/δ phase boundary were observed by SEM (Fig. 2c). 2.2 Analysis of surface microstructure

The Vickers microhardness of the primary surface and cross-section of the

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specimens after different surface treatments was obtained using a microhardness tester

(HXD-1000TMB) with a load of 100 g and a hold time of 15 s. For each specimen, the

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microhardness is measured for at least five times. The cross-section specimens were exacted by wire electrical discharge machining, then mechanically ground by

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waterproof abrasive paper up to 2000 grit, and finally polished using 1 μm diamond paste.

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The surface morphology was characterized using a field emission SEM (FEI Quanta 650). A 3D Profiler (VX-X100K) was used to analyze the surface roughness of the specimens after different surface treatments. Each surface treatment was measured

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three times with an analyzed area of 2 mm × 1.5 mm. Ra, the arithmetic average of three profiles was chosen to represent the surface roughness.

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The primary surface after the different surface treatments was analyzed by XPS

(ESCALAB 250). The depth-profile XPS analyses were performed by using argon ion beam sputtering an area of 2 mm diameter referred to Ta2O5 layer at a rate of 0.1 nm/s and a target current density of 2 A/cm2. A focusing X-ray monochromator with a 0.5 mm spot size was used to obtain high-resolution spectra. Spectra of Cr 2p3/2, Fe 2p3/2, Ni 2p3/2, O 1s and C 1s were systematically recorded. Prior to the peak fitting of the 6

XPS spectrum, the binding energies were calibrated using the C1s peak at 285 eV. Quantification of the species in the primary surface were performed via XPSpeak4.1 peak fitting software using Gaussian–Lorentzian peak shapes and a Shirley background. Chemical states of Cr, Ni, and Fe were analysed by normalizing values of each element peak area. Binding energies of the compound for fitting each element spectra were obtained from the literature [40]. Cross-section microstructure and residual strain after the different surface treatments were analyzed using an EBSD detector attachment in the SEM equipped

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with a camera and TSL software. The specimens used for EBSD analysis were ground using SiC papers up to 2000 grit, then mechanically polished by diamond paste to 1 μm,

and finally polished with 40 nm colloidal silica slurry. The EBSD analyses of the cross-

section of the specimens were performed with a step size of 2 μm at a voltage of 25

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keV. Local strain distribution in the weld metal was assessed using kernel average

misorientation (KAM), which has a linear relationship with the degree of strain [41].

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KAM is a local misorientation defined as an average misorientation of a point with respect to all of its neighbors in a grain. For a given point, the average misorientation

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of that point was calculated with a criterion that misorientations exceeding a tolerance value (5º here) were excluded from the calculation to avoid grain boundary effects [42].

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After the EBSD analyses, an Anton Paar UNHT nanoindenter with a Berkovich diamond indenter was used to analyze the nanoindentation hardness of the δ ferrite and γ austenite phases. The nanoindentation tests were conducted on the cross-section of

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the specimens with a distance of about 20 μm to the specimen surface. Both the loading rate and the unloading rate were 0.1 mN/s with a load of 3 mN and a pause time of 5 s.

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At least five measurements were performed for a reproducibility of the measurement. Grain boundary microstructure of the bulk metal, the underlying deformed layer

and the nanocrystalline layer of the milled specimen were analyzed using a JEOL 2100 TEM equipped with an energy dispersive X-ray spectroscopy detector (EDX). For the TEM analysis of the bulk metal, specimens with an original dimension of 10 mm × 10 mm × 0.2 mm were cut by wire electrical discharge machining. Then they were ground 7

using SiC papers up to 2000 grit to reduce the thickness below 50 μm. A hole drill bit with a 3 mm inner diameter was used to create discs through the thinned specimens. These discs were then thinned to electron transparency by a Struers Tenupol-5 Twin-jet electropolishing system, using a solution of 5 vol.% perchloric acid of 70 wt.% concentration and 95 vol.% ethanol of 99.7 wt.% concentration at 25 V and −30–−25 °C. To analyze the microstructure of the underlying deformed layer and the nanocrystalline layer, specimens were fabricated using a dual-beam focused ion beam (FIB) instrument (Helios 600i, FEI).

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2.3 SCC test A refreshed loop equipped with a 10 L autoclave made of 316L stainless steel was

used for the SCC test in the primary PWR environment at 320 °C and 13 MPa with a

flow rate at 5 L/h. The primary PWR environment was prepared by high-purity water

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with 1.2 g/L of B as H3BO3 and 2.2 mg/L of Li as LiOH. Before adding the chemicals

into the water tank, the water in the loop was circulated and purified with a high-

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efficiency ion-exchanger to a conductivity of 0.067 μS/cm.

The bent beam specimens were bent and fixed by an upper fixture with a concave

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surface and a lower fixture with a convex surface (Fig. 1d). Details of the fixture and testing process were reported previously [27]. Compared to the crevice bent beam tests

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used in literature [43, 44], it may take a longer time for SCC initiation at the same tensile strain during the bent beam tests due to the free surface. To accelerate crack initiation, a maximum strain of 2% at the specimen surface was applied in the tests.

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For each surface treatment, three specimens were employed in SCC tests. Dissolved hydrogen in the SCC test was controlled by bubbling H 2 into the water tank

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until equilibration occurred. The concentration of dissolved hydrogen and oxygen in the influent water was 2.6 mg/L and <5 μg/L, respectively. Prior to heating the autoclave, the specimens were exposed to the water at room temperature for 6 h to stabilize the environmental condition. Then the autoclave was heated up to 320 C and maintained at this temperature for 2500 h without interruption. 2.4 Analyses of the SCC behavior 8

Following the SCC tests, each specimen was removed from the autoclave for surface observation with an area of 20 mm × 10 mm by SEM. They were then sliced into pieces along the longitudinal direction to observing the cracks in cross-section. Since the cracks only initiated on convex surface of the specimen, three regions adjacent to the convex edge with a length of 40 mm were selected for counting the number of cracks and measuring the maximum and average crack depth for each surface treatment. The crack concentration was defined as the number of cracks per unit length, as shown in the equation:

 = n/L

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(1)

where  is the crack concentration, n is the total number of the observed cracks, L is

the total length of the observed region. The unit of crack concentration is cm−1. It is worth noting that only the cracks deeper than 1 μm were selected for the statistics. The

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SCC susceptibility after the different surface treatments was primarily evaluated by the maximum and average crack depth.

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EBSD analysis was also conducted on the cross-section of the milled specimen to reveal the crack propagation path and residual strain distribution. In addition, the crack

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tips in the nanocrystalline layer on the milled specimen were evaluated by TEM. The TEM samples of crack tips were prepared by FIB.

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It should be mentioned that the SCC behavior of weld cladding may not be broadly relevant to structural welds due to the differences in residual stresses. 3. Results

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3.1 Characterization of the specimen after surface treatments 3.1.1 Surface morphology, roughness and microhardness

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Surface morphology after the different surface treatments was analyzed by SEM

and 3D Profiler as shown in Fig. 3 and Fig. 4, respectively. The slurry polished specimen had a very smooth surface while peaks and valleys can be observed on the surface of ground specimen (Fig. 3a and b, Fig. 4a and b). On the milled specimen, the machining marks were observed on its surface (Fig. 3c and 4c). Island-shaped δ ferrite phase appeared on the surface of the electropolished specimen (Fig. 3d and 4d). In 9

addition, the phase boundaries were slightly etched after electropolishing. The surface microhardness and roughness after the different surface treatments are listed in Table 2. The slurry polished and electropolished specimens had a lower microhardness than the ground and milled specimens. On the machining marks of the milled specimen, the microhardness of the peak was higher than in the valley. The diagonal of the rhombic indentation was about 20 μm (Fig S1). Therefore, according to the calculated depth of indentation (~4 μm) and the thickness of nanocrystalline layer (see Fig. 5), only the nanocrystalline layer of the milled specimen was measured. The

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surface roughness of the slurry polished specimen was the lowest among the different surface treatments, whereas the ground specimen showed the highest surface roughness. In addition, the surface roughness of the milled specimen was half that of the ground specimen.

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3.1.2 Microstructure, microhardness and residual strain distribution, and composition depth profile in cross-sections

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Fig. 5 shows the backscattered electron images of the cross-section of the milled specimen. A 5–10 μm thick nanocrystalline layer was observed on the specimen surface

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(Fig. 5a and b), which is a combined effect of severe plastic shear deformation and high temperature [45]. The nanocrystalline layer mainly consisted of nano-sized austenite

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and elongated δ ferrite phases (Fig. 5b and c). The δ ferrite phase was broken up into nano-sized grain after severe plastic deformation, which was confirmed by Transmission Kikuchi Diffraction [15]. In addition, there was a deformed layer

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beneath the nanocrystalline layer (Fig. 5d). The slip lines in the deformed layer tended not to go through the δ ferrite phase, but to terminate at the γ/δ phase boundary. The

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near-surface microstructure of the ground specimen (Fig S2) was similar to the milled specimen, which is consistent with the literature [20]. Fig. 6 shows the load-depth curves of δ ferrite and austenite phases with a distance

of about 20 μm to the specimen surface after different surface treatments. The nanoindentation hardness of both the δ ferrite and γ austenite phases decreased successively from the milled specimen, the ground specimen to the slurry 10

polished/electropolished specimen, indicating the hardening of both δ ferrite and austenite phases during surface deformation. Further, the δ ferrite phase always showed a higher hardness than the austenite phase. The KAM and microhardness distribution in the cross-section after the different surface treatments are shown in Fig. 7. The surfaces of milled and ground specimens show higher residual strains than the slurry polished and electropolished specimens. This is consistent with the microhardness distribution (Fig. 7f) and nanoindentation hardness (Fig. 6). Increasing the distance from the specimen surface led to decrease of

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KAM value and microhardness (Fig. 7e and f). The KAM analyses and microhardness distribution revealed the thickness of the deformed layer on the surface of the milled

specimen was about 100 μm, which was thicker than on the ground specimen. It should

be mentioned that, since the grain size of the nanocrystalline layer was very small, the

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KAM distribution in the nanocrystalline layer cannot be detected.

Composition depth profiles of Cr, Ni, and Fe with different valence states after the

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different surface treatments before SCC tests are shown in Fig. 8. A thin film consisting of Cr hydroxide, Cr oxide, and Fe oxide was observed on all specimens. The surface

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layer of the film showed enrichment of Cr. The concentration of Cr element on the surface of the milled specimen was about 35 at.%, which was lower than for other

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specimens. Further, the Cr hydroxide showed a large percentage than the Cr oxide in the surface layer of the film. It should be mentioned that the elemental Fe and Ni measured by the XPS analysis after sputtering for 10 s was likely because the thickness

46].

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of film remained on the surface is thinner than the analysis depth of XPS (~5 nm) [27,

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TEM observations and analyses of the bulk metal, underlying deformed layer and

nanocrystalline layer of the milled specimen are shown in Fig. 9. Carbides of 100–200 nm existed as a semi-continuous distribution along the γ/δ phase boundary that were observed in both the bulk metal and underlying deformed layer (Fig. 9a and b). As reported in literature [47], the carbides in 308L weld metal are M23C6 type. The density of dislocations in the bulk metal was much lower than in the underlying deformed layer. 11

The nanocrystalline layer mainly consisted of nano-sized austenite grains, as confirmed by the selected-area diffraction (Fig. 9c). The morphology of nano-sized austenite grains was not equiaxed and was similar to the ferrite grains (Fig. 9c and d). The grain size of nano-sized austenite and ferrite along the milling direction was much larger than that along the depth direction. In addition, numerous dislocations and a few fractured carbides with a size of about 30 nm were observed in the nanocrystalline layer (Fig. 9d). The EDX mappings and line scans shown in Fig. 9e and f revealed slight depletion of Cr (~22 at.%) at the γ/δ and δ/M23C6 boundary. The depletion of Cr also resulted in

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enrichment of Ni at the phase boundary (Fig. 9g). In addition, the composition distribution of the phase boundary in the nanocrystalline layer was similar to the bulk

metal (Fig. 9g and h), i.e., the Cr content of the δ ferrite phase decreased from the center to the phase boundary.

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3.2 SCC behavior after the different surface treatments in the primary PWR environment

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3.2.1. Observation of cracking on surface

SEM observations of the surface of the specimens with different surface treatments

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after SCC tests are shown in Fig. 10. No cracks, and only discrete oxide particles with a size of less than 1 μm, were observed on the slurry polished specimen (Fig. 10a).

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However, there were a few irregular cracks and oxide particles on the surface of ground specimen (Fig. 10b). With regard to the milled specimen, intergranular cracks initiated in the valley and propagated along the machining grooves (Fig. 10c). Further, it was

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found that the size and density of oxide particles formed on the milled specimen were larger than on the other specimens, as shown in the inset figure. SCC initiation at the

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γ/δ phase boundary was also observed on the electropolished specimen (Fig. 10d). 3.2.2. Observation and analyses of cracking in cross-sections SEM images of the cracking in cross-sections after SCC tests are shown in Fig. 11.

It was confirmed that no cracks initiated on the slurry polished specimen (Fig. 11a). However, a few shallow cracks that were always parallel to the surface were observed on the ground specimen (Fig. 11b). Representative SCC cracks on the milled specimen 12

are shown in Fig. 11c-e. The cracks were much longer than on the ground specimen. The elongated δ ferrite phase in the nanocrystalline layer were a preferential path for SCC initiation, as is evident by the filamentous oxidation along the elongated δ ferrite phase. Uniform oxide film with a thickness of ~1 μm was also observed on the surface of milled specimen (Fig. 11e). The uniform oxide film on the surface of milled specimen was much thicker than on the electropolished specimen (Fig. 11f). This was related to the fact that a high density of grain boundaries in the nanocrystalline layer promoted faster oxidation of the milled specimen compared to the electropolished

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specimen [20]. The shallow cracks on the electropolished specimen were along the γ/δ phase boundary (Fig. 11f).

The concentration, maximum and average depth of the cracks on the cross-section of the specimen after SCC tests are summarized in Fig. 12. The maximum and average

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depth of the crack observed on the milled specimen were much bigger than other

specimens, indicating the highest SCC susceptibility. In addition, while the

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electropolished specimen shows the highest crack concentration, the average crack depth was only about 3 μm.

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3.2.3 Observation and analysis of crack growth of the milled specimen SEM observation and EBSD analysis of cracking on cross-sections of the milled

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specimen in the primary PWR environment are shown in Fig. 13. A crack deeper than 100 μm was observed on the milled specimen (Fig. 13a). SCC propagated along the γ/δ phase boundary while apparent cessation of crack growth at the γ/γ grain boundary was

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observed (Fig. 13b). The phase boundary that intersected with deformation bands was the preferred site for SCC growth (Fig. 13c). This suggests the phase boundary was

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subjected a higher deformation and was more susceptible to SCC. In addition, macrocavities were observed at the metal-oxide interface, which may due to the different lattice parameters of the metal matrix compared to the oxides and the high density of vacancies caused by deformation [48, 49] (Fig. 13d). As shown in Fig. 13e, the inverse pole figure revealed the grain orientation of the γ austenite phase was different to the δ ferrite phase. In addition, there was strain concentration adjacent to boundaries in the 13

SCC area and near the γ/δ phase boundary according to the KAM map. 3.2.4 TEM observation and analysis of cracking in the nanocrystalline layer Since grinding and milling produced a similar nanocrystalline layer, TEM specimens extracted from the milled specimen were analyzed to investigated the effect of nanocrystalline layer on SCC. Morphology and composition profiles of a crack that intersects with the elongated δ ferrite phase in the nanocrystalline layer are shown in Fig. 14. The bright-field TEM image shows the elongated δ ferrite phase was completely oxidized and no nano-pores were observed in the oxide (Fig. 14a and b).

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The selected-area diffraction pattern and the EDX line scan across the oxidized δ ferrite phase confirmed that the center of the oxide was maghemite/magnetite (Fig. 14c and d). Preferential oxidation was observed along the γ/δ phase boundary. The EDX line

scans show no enrichment of Cr in the oxide, suggesting a low protectiveness of the

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oxide adjacent to the δ ferrite phase (Fig. 14e and f).

Fig. 15 shows the morphology and composition analyses of cracks that intersect

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with carbides, and the austenite phase in the nanocrystalline layer. An increase in Fe content but decrease in O content was observed from the oxidized δ ferrite phase

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towards the carbide, suggesting that diffusion of oxygen was restrained by the high Cr content in the carbide (Fig. 15c). The beneficial effect of carbides on oxidation is

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consistent with the literature [48]. In addition, the oxide adjacent to the austenite phase has a higher Cr content but lower Fe content, suggesting that the oxide was protective (Fig. 15e). The high Cr content in the oxide was due to Cr diffusion from the adjacent

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austenite matrix. 4. Discussion

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4.1 Effect of electropolishing on SCC of 308L weld metal in the primary PWR environment

The SEM observations and analyses shown in Figs. 10-12 reveal a higher SCC

susceptibility of the electropolished specimen than the slurry polished specimen. Since both the colloidal silica slurry polishing and electropolishing produce a similar surface residual strain and roughness, the preferential initiation of SCC on the electropolished 14

specimen is likely due to the composition change of the surface. During electropolishing, enhanced dissolution of alloying elements could occur at the γ/δ phase boundary resulting in etched phase boundaries (Fig. 3). As revealed by the XPS analysis, more chromium hydroxides formed on the surface of electropolished specimens than on slurry polished specimens. Since the Cr hydroxide was less protective than the Cr oxide, O could more easily penetrate the etched γ/δ phase boundaries [27], resulting in SCC initiation of 308L weld metal. It is worthwhile to mention that a higher corrosion rate of 308L weld metal after electropolishing than slurry polishing was also observed

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in literature [37]. There is conflicting data on the effect of electropolishing on SCC of materials in

primary PWR environments. Recent research on Alloy 182 revealed the electropolishing promoted the formation of a protective oxide at grain boundaries,

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which could mitigate the intergranular corrosion and reduce SCC susceptibility [27]. The beneficial effect of electropolishing on suppressing SCC were also observed in 304

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stainless steel [25]. The different effect of electropolishing on SCC between the 308L weld metal and other alloys is likely related to the oxidation behavior of the etched

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boundary. Further work at relatively short exposure time is needed to better understand this aspect.

environment.

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4.2 Effect of nanocrystalline layer on SCC of 308L weld metal in the primary PWR

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The nanocrystalline layer may increase the SCC initiation resistance of 304L/316L stainless steel, because of more uniform oxidation in the nanocrystalline layer [13-17].

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However, in contrast to these observations, the results in the present work showed that the nanocrystalline layer increased the SCC initiation susceptibility. Microstructure and residual stress/strain are two dominant material factors in SCC initiation. Although high residual stress exists in the nanocrystalline layer of the stainless steel after milling [16], the effect of residual stress in the nanocrystalline layer is unlikely to be important because the surface stress of the material leveled off to similar values after bending [13, 14, 16]. Further, the SEM and TEM observation shown in Figs. 11 and 13 revealed that 15

the elongated δ ferrite phase in the nanocrystalline layer was most likely oxidized in the primary PWR environment. This suggests the elongated δ ferrite phase is a preferential path for SCC initiation. However, it has been extensively reported that δ ferrite phase has a beneficial effect on SCC of stainless steels in high temperature water [15, 33, 39, 49, 50]. This discrepancy may arise from the difference in the concentration of dislocations and the size of δ ferrite phase. According to the literature [15, 33, 49, 50], SCC susceptibility of stainless steel was decreased by the δ ferrite phase, likely due to the inhibition of oxidation and sliding

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of grain boundary. When cracks intersect with the δ ferrite phase, it acts as a Cr source to facilitate the formation of Cr-enriched oxides (Cr2O3 and Fe-Cr spinel) at phase

boundaries, due to the high diffusion rate and content of Cr in δ ferrite phase [49, 50].

The Cr-enriched oxides could impede inward diffusion of oxygen at the phase boundary,

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and decrease the SCC susceptibility [49, 50]. The grain boundary sliding is one of the

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key steps for SCC in high hydrogenated temperature. Carbide nucleation and growth occur first at the incoherent ferrite/austenite phase boundary during welding and post-

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weld heat treatment (Figs. 2 and 9). The carbides could act as a barrier to prevent grain boundary sliding and increased the SCC resistance of stainless steel in high temperature hydrogenated water [51-53]. In addition, for the dispersed island-shaped ferrite phases,

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another mechanical effect is the diversion of the crack path, which increases the total crack length and decreases the effective stress at each crack tip [49, 50]. Thus, the SCC

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growth of the stainless steels is inhibited by the presence of δ ferrite phase. In the current study, the detrimental role of δ ferrite phase in SCC of 308L weld

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metal after machining in the simulated primary PWR environment has been drawn, as summarized in Fig. 16. When exposed in the primary PWR environment, a uniform oxide film formed on the surface due to the high density of grain boundaries in the nanocrystalline layer (Fig. 16b). Further, since the higher density of dislocations adjacent to the γ/δ phase boundary could act as short circuits for diffusion of oxygen and metallic elements, and thus resulted in preferential corrosion along phase boundaries in the nanocrystalline layer [48, 54-57]. Due to the higher diffusion rate of 16

Cr in the δ ferrite phase than the austenite phase, outward diffusion of Cr occurred from the δ ferrite matrix towards the phase boundary [58, 59]. Nevertheless, the nano-sized δ ferrite phase could not provide sufficient Cr for the formation of a Cr-enriched oxide, but led to a Cr-depleted oxide in the center of δ ferrite phase (Fig. 14). Thus, the filamentous δ ferrite phase was severely oxidized by the inward diffusion of O (Fig. 16b). After an amount of time, cavities formed at the metal-oxide interface, which acted as short-circuits for water ingress [48]. Therefore, fast oxidation of δ ferrite phase could continue, leading to accelerated SCC initiation in the nanocrystalline layer. Following

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fast oxidation of elongated δ ferrite phase in the nanocrystalline layer, those cracks can propagate intergranularly into the underlying deformed layer along the phase boundary,

due to the high dislocation density and the lath-like morphology of δ ferrite phase (Fig. 16c). In contrast to the corrosion of elongated δ ferrite phase, Cr-enriched oxide could

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therefore oxidation will be reduced (Fig. 15).

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form at the interface once the oxidation meets a carbide or the austenite phase, and

5. Conclusion

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This study investigated the effect of surface treatments on microstructure and stress corrosion cracking behavior of 308L weld metal exacted from the cladding on SA508 low alloy steel in a primary PWR environment. The following conclusions were drawn:

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(1) SCC was observed on the milled, mechanically ground and electropolished surfaces, but not on the colloidal silica slurry polished surface.

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(2) The change in composition and etched phase boundary at the surface after electropolishing promotes SCC initiation of 308L weld metal.

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(3) Milling and grinding produce a nanocrystalline layer and an underlying

deformed layer on the surface of 308L weld metal. In the nanocrystalline layer of the milled surface, the length of nano-sized austenite and ferrite along milling direction was much greater than along the depth direction. (4) The elongated δ ferrite phase in the nanocrystalline layer has a detrimental effect on SCC initiation of 308L weld metal. This is because the synergistic effect of high concentration of dislocations and the nano size of δ ferrite phase did not provide 17

sufficient Cr for the formation of Cr-enriched oxides, but lead to a lower Cr in the center of the δ ferrite phase.

Author statement

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Lijin Dong: Investigation, Validation, Data Curation, Writing - Original Draft, Writing - Review & Editing. Xiaolong Zhang: Investigation, Visualization Yaolei Han: Investigation Qunjia Peng: Supervision, Writing - Review & Editing Ping Deng: Investigation Shuliang Wang: Writing - Review & Editing Declaration of interests

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The authors declare that they have no known competing financial interests or personal

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Data availability

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relationships that could have appeared to influence the work reported in this paper.

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

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time as the data also forms part of an ongoing study

Acknowledgements

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This work is supported by Youth Scientific and Innovation Research Team for Advanced Surface Functional Materials (No. 2018CXTD06), Open Research Fund

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from State Key Laboratory of Metal Material for Marine Equipment and Application (SKLMEA-K201912), Scientific Research Starting Project of SWPU (No. 2019QHZ012), Young Scholars Development Found of Southwest Petroleum University (No. 201899010040), and the China Postdoctoral Science Foundation (2018M642319). The authors are also grateful to Ms. X. Liang and Dr. Y.F. Li for their support to perform TEM analyses. 18

Refences: [1] A. Turnbull, K. Mingard, J.D. Lord, B. Roebuck, D.R. Tice, K.J. Mottershead, N.D. Fairweather, A.K. Bradbury, Sensitivity of stress corrosion cracking of stainless steel to surface machining and grinding procedure, Corr. Sci. 53 (2011) 3398-3415. https://doi.org/10.1016/j.corsci.2011.06.020 [2] P.L. Andresen. Stress corrosion cracking of current structural materials in commercial nuclear power plants. Corrosion, 69 (2013) 1024-1038.

ro of

[3] D.R. Tice, V. Addepalli, K.J. Mottershead, M.G. Burke, F. Scenini, S. Lozano-Perez, G. Pimentel, Microstructural effects on stress corrosion initiation in austenitic stainless steel in PWR environments, in: Proceedings of the 18th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Portland, Oregon, USA, August 13–17, 2017, pp. 775792.

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[4] R. Obata, M. Koshiishi, H. Anzai, K. Nakade, S. Ooki, S. Suzuki, Correlation between oxide film and stress corrosion cracking susceptibility of surface cold worked L-grade stainless steels, in: Proceedings of the 14th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Virginia Beach, Virginia, August 23-27, 2009, pp. 622-634.

lP

re

[5] X.Y. Zhong, S.C. Bali, T. Shoji, Accelerated test for evaluation of intergranular stress corrosion cracking initiation characteristics of non-sensitized 316 austenitic stainless steel in simulated pressure water reactor environment, Corr. Sci. 115 (2017) 106-117. https://doi.org/10.1016/j.corsci.2016.11.019

ur

na

[6] G. Pimentel, D. Tice, V. Addepalli, K. Mottershead, M. Burke, F. Scenini, J. Lindsay, Y. Wang, S. Lozano-Perez, High-resolution characterisation of austenitic stainless steel in PWR environments: effect of strain and surface finish on crack initiation and propagation, in: Proceedings of the 18th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Portland, Oregon, USA, August 13–17, 2017, pp. 829-847.

Jo

[7] L.J. Dong, Q.J. Peng, E.-H. Han, W. Ke, L. Wang, Stress corrosion cracking in the heat affected zone of a stainless steel 308L-316L weld joint in primary water, Corr. Sci. 107 (2016) 172-181. https://doi.org/10.1016/j.corsci.2016.02.030 [8] L.J. Dong, Q.J. Peng, E.-H. Han, W. Ke, L. Wang, Microstructure and intergranular stress corrosion cracking susceptibility of a SA508-52M-316L dissimilar metal weld joint in primary water, J. Mater. Sci. Technol. 34 (2018) 1281-1292. https://doi.org/10.1016/j.jmst.2017.11.051 [9] R.L. Zhu, J.Q. Wang, L.T. Zhang, Z.M. Zhang, E.-H. Han, Stress corrosion cracking of 316L HAZ for 316L stainless steel/Inconel 52M dissimilar metal weld joint in simulated primary water, Corr. Sci. 112 (2016) 373-384. 19

https://doi.org/10.1016/j.corsci.2016.07.031 [10] M.J. Olszta, L.E. Thomas, K. Asano, S. Ooki, S.M. Bruemmer, Crack initiation precursors originating from surface grinding, in: Proceedings of the 14th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Virginia Beach, Virginia, August 23-27, 2009, pp. 549-561. [11] J. Kaneda, H. Tamako, R. Ishibashi, H. Hato, M. Miyagawa, N. Yamashita, Effects of surface treatments on microstructure, hardness and residual stress in type 316L stainless steel, in: Proceedings of the 14th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Virginia Beach, Virginia, August 23-27, 2009, pp. 791-802.

ro of

[12] Z.Q. Zhai, M.B. Toloczko, M.J. Olszta, S.M. Bruemmer, Stress corrosion crack initiation of alloy 600 in PWR primary water, Corr. Sci. 123 (2017) 76-87. https://doi.org/10.1016/j.corsci.2017.04.013

-p

[13] L.T. Chang, M.G. Burke, F. Scenini, Stress corrosion crack initiation in machined type 316L austenitic stainless steel in simulated pressurized water reactor primary water, Corr. Sci. 138 (2018) 54-65. https://doi.org/10.1016/j.corsci.2018.04.003

lP

re

[14] L.T. Chang, J. Duff, M.G. Burke, F. Scenini, SCC initiation in the machined austenitic stainless steel 316L in simulated PWR primary water, in: Proceedings of the 18th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Portland, Oregon, USA, August 13–17, 2017, pp. 811-827.

na

[15] L.T. Chang, M.G. Burke, F. Scenini, Understanding the effect of surface finish on stress corrosion crack initiation in warm-forged stainless steel 304L in high-temperature water, Scripta. Mater. 164 (2019) 1-5. https://doi.org/10.1016/j.scriptamat.2019.01.032

Jo

ur

[16] L.T. Chang, L. Volpe, Y.L. Wang, M.G. Burke, A. Maurotto, D. Tice, S. Lozano-Perez, F. Scenini, Effect of machining on stress corrosion crack initiation in warm-forged type 304L stainless steel in high temperature water, Acta Mater. 165 (2019) 203-214. https://doi.org/10.1016/j.actamat.2018.11.046 [17] F. Scenini, J. Lindsay, L.T. Chang, Y.L. Wang, M.G. Burke, S. Lozano-Perez, G. Pimentel, D. Tice, K. Mottershead, V. Addepalli, Oxidation and SCC initiation studies of type 304L SS in PWR primary water, in: Proceedings of the 18th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Portland, Oregon, USA, August 13–17, 2017, pp. 793810. [18] Y. Katayama, M. Tsubota, Y. Saito, Effect of crystal orientation on the stress corrosion cracking of directionally solidified L-grade stainless steels, in: Proceedings of the 13th Conference on 20

Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Whistler, British Columbia, August 19 - 23, 2007, CD–ROM. [19] L. Wickström, K. Mingard, G. Hinds, A. Turnbull, Microcrack clustering in stress corrosion cracking of 22Cr and 25Cr duplex stainless steels, Corr. Sci. 109 (2016) 86-93. https://doi.org/10.1016/j.corsci.2016.03.024 [20] S.Y. Wang, Y.J. Hu, K. Fang, W.Q. Zhang, X.L. Wang, Effect of surface machining on the corrosion behaviour of 316 austenitic stainless steel in simulated PWR water, Corr. Sci. 126 (2017) 104-120. https://doi.org/10.1016/j.corsci.2017.06.019

ro of

[21] F. Scenini, R.C. Newman, R.A. Cottis, R.J. Jacko, Effect of surface preparation on intergranular stress corrosion cracking of Alloy 600 in hydrogenated steam, Corrosion, 64 (2008) 824-835. https://doi.org/10.5006/1.3279916

-p

[22] F. Scenini, R.C. Newman, R.A. Cottis, R.J. Jacko, Alloy oxidation studies related to PWSCC, in: Proceedings of the 12th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Salt Lake City, Utah, August 14–18, 2005, pp. 891-902.

re

[23] M. Payet, L. Marchetti, M. Tabarant, J.-P. Chevalier, Corrosion mechanism of a Ni-based alloy in supercritical water: Impact of surface plastic deformation, Corr. Sci. 100 (2015) 47-56. https://doi.org/10.1016/j.corsci.2015.06.032

na

lP

[24] S. Habibzadeh, L. Li, D. Shum-Tim, E.C. Davis, S. Omanovic, Electrochemical polishing as a 316L stainless steel surface treatment method: Towards the improvement of biocompatibility, Corr. Sci. 87 (2014) 89-100. https://doi.org/10.1016/j.corsci.2014.06.010

ur

[25] H. D. Solomon, Transgranular, granulated, and intergranular stress corrosion cracking in AISI 304 SS, Corrosion, 40 (1984) 493-506. https://doi.org/10.5006/1.3577922

Jo

[26] P. L. Andresen, Crack initiation in cert tests on type 304 stainless steel in pure water, Corrosion, 38 (1982) 53-59. https://doi.org/10.5006/1.3577319 [27] Y. L. Han, E.-H. Han, Q.J. Peng, W. Ke, Effects of electropolishing on corrosion and stress corrosion cracking of Alloy 182 in high temperature water, Corr. Sci. 121 (2017) 1-10. https://doi.org/10.1016/j.corsci.2017.03.004 [28] H. Solomon, M. Povich, T. Devine, Slow strain-rate testing in high temperature water, in: G. Ugiansky and J. Payer, (Eds.), Stress Corrosion Cracking—The Slow Strain-Rate Technique, ASTM International, West Conshohocken, 1979, 132-148. 21

https://doi.org/10.1520/STP38113S [29] P.L. Andresen, Stress corrosion cracking (SCC) of austenitic stainless steels in high temperature light water reactor (LWR) environments, in: P. G. Tipping (Eds.), Understanding and Mitigating Ageing in Nuclear Power Plants, Woodhead Publishing, 2010, pp. 236-307. https://doi.org/10.1533/9781845699956.2.236. [30] G. Sui, E. A. Charles, J. Congleton, The effect of delta-ferrite content on the stress corrosion cracking of austenitic stainless steels in a sulphate solution, Corr. Sci. 38 (1996) 687-703. https://doi.org/10.1016/0010-938X(96)00159-X

ro of

[31] T. M. Devine, B. J. Drummond, Use of accelerated intergranular corrosion tests and pitting corrosion tests to detect sensitization and susceptibility to intergranular stress corrosion cracking in high temperature water of duplex 308 stainless steel, Corrosion, 37 (1981) 104-115. https://doi.org/10.5006/1.3593843

-p

[32] L.J. Dong, E.-H. Han, Q.J. Peng, W. Ke, L. Wang, Environmentally assisted crack growth in 308L stainless steel weld metal in simulated primary water, Corr. Sci. 117 (2017) 1-10. https://doi.org/10.1016/j.corsci.2016.12.011

re

[33] H. Abe, Y. Watanabe, Role of δ-ferrite in stress corrosion cracking retardation near fusion boundary of 316NG welds, J. Nucl. Mater. 424 (2012) 57-61. https://doi.org/10.1016/j.jnucmat.2012.02.006

na

lP

[34] T. Lucas, A. Forsstrom, T. Saukkonen, R. Ballinger, H. Hanninen, Effects of thermal aging on material properties, stress corrosion cracking, and fracture toughness of AISI 316L weld metal, Metall. Mater. Trans. A 47 (2016) 3956-3970. https://doi.org/10.1007/s11661-016-3584-6

ur

[35] Metallographic investigation on the cladding failure in the pressure vessel of a BWR, Nucl. Eng. Des. 16 (1971) 205-222. https://doi.org/10.1016/0029-5493(71)90010-0

Jo

[36] H.Q. Xu, S. Fyfitch, Laboratory investigation of PWSCC of CRDM nozzle 3 and its J-groove weld on the Davis-Besse reactor vessel head, in: Proceedings of the 12th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Salt Lake City, Utah, August 14–18, 2005, pp. 833-842. [37] C. Ma, E.-H. Han, Q.J. Peng, W. Ke, Effect of polishing process on corrosion behavior of 308L stainless steel in high temperature water, Appl. Surf. Sci. 442 (2018) 423-436. https://doi.org/10.1016/j.apsusc.2017.12.190 [38] H.L. Ming, Z.M. Zhang, J.Q. Wang, R.L. Zhu, J. Ding, J.Z. Wang, E.-H. Han, W. Ke, Effect of surface state on the oxidation behavior of welded 308L in simulated nominal primary water of PWR, 22

Appl. Surf. Sci. 337 (2015) 81-89. https://doi.org/10.1016/j.apsusc.2015.02.066 [39] K. Mukahiwa, G. Bertali, M.G. Burke, J. Duff, F. Scenini, The beneficial effect of surface carbon coating on stress corrosion cracking of Type 304 austenitic stainless steels in high temperature water, Scripta. Mater. 158 (2019) 77-82. https://doi.org/10.1016/j.scriptamat.2018.08.033 [40] Y.L. Han, J.N. Mei, Q.J. Peng, E.-H. Han, W. Ke, Effect of electropolishing on corrosion of nuclear grade 316L stainless steel in deaerated high temperature water, Corr. Sci. 112 (2016) 625634. https://doi.org/10.1016/j.corsci.2016.09.002

ro of

[41] M. Kamaya, Measurement of local plastic strain distribution of stainless steel by electron backscatter diffraction, Mater. Charact. 60 (2009) 125-132. https://doi.org/10.1016/j.matchar.2008.07.010

-p

[42] D.W. Gross, K. Nygren, G.J. Pataky, J. Kacher, H. Sehitoglu, I.M. Robertson, The evolved microstructure ahead of an arrested fatigue crack in Haynes 230, Acta Mater. 61 (2013) 5768-5778. https://doi.org/10.1016/j.actamat.2013.06.020

lP

re

[43] Y. Katayama1, M. Tsubota, Y. Saito, Effect of crystal orientation on the stress corrosion cracking of directionally solidified L-grade stainless steels, in: 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, TMS, 2007, CD–ROM.

na

[44] H. P. Offer, R.M. Horn, A.Q. Chan, Assessment of the mitigation of SCC by surface stress and material improvements, in: 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, TMS, 2007, CD–ROM.

Jo

ur

[45] M. Imran, P.T. Mativenga, A. Gholinia, P.J. Withers, Comparison of tool wear mechanisms and surface integrity for dry and wet micro-drilling of nickel-base superalloys, Int. J. Mach. Tool. Manu. 76 (2014) 49-60. https://doi.org/10.1016/j.ijmachtools.2013.10.002 [46] W.D. Zhang, E. Bunte, F. Ruske, D. Köhl, A. Besmehn, J. Worbs, H. Siekmann, J. Kirchhoff, A. Gordijn, J. Hüpkes, As-grown textured zinc oxide films by ion beam treatment and magnetron sputtering, Thin solid films, 520 (2012) 4208-4213. https://doi.org/10.1016/j.tsf.2011.04.098 [47] X.D. Lin, Q.J. Peng, E.-H. Han, W. Ke, C. Sun, Z.J. Jiao, Irradiation-induced segregation at phase boundaries in austenitic stainless steel weld metal, Scripta. Mater. 149 (2018) 11-15. https://doi.org/10.1016/j.scriptamat.2018.01.033 23

[48] Z. Shen, J. Dohr, S. Lozano-Perez, The effects of intergranular carbides on the grain boundary oxidation and cracking in a cold-worked Alloy 600, Corr. Sci. 115 (2019) 209-216. https://doi.org/10.1016/j.corsci.2019.05.008 [49] J.M. Wang, H.Z. Su, K. Chen, D.H. Du, L.F. Zhang, Z. Shen, Effect of δ-ferrite on the stress corrosion cracking behavior of 321 stainless steel, Corr. Sci. 158 (2019) 108079 https://doi.org/10.1016/j.corsci.2019.07.005 [50] Y. Watanabe, H. Abe, Role of delta–ferrite in SCC retardation near fusion boundary of 316LC welds, in: Minutes of the 2008 annual meeting of the International Cooperative Group on Environmentally Assisted Cracking of water reactor materials, Bastad, Sweden, April 20-25, 2008, CD–ROM.

ro of

[51] K. Arioka, T. Yamada, T. Terachi, R.W. Staehle, Intergranular stress corrosion cracking behavior of austenitic stainless steels in hydrogenated high-temperature water, Corrosion, 62 (2006) 74-83. https://doi.org/10.5006/1.3278254

re

-p

[52] W.J. Kuang, G.S. Was, The effects of grain boundary carbide density and strain rate on the stress corrosion cracking behavior of cold rolled Alloy 690, Corr. Sci. 97 (2015) 107-114. https://doi.org/10.1016/j.corsci.2015.04.020

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[53] L.J. Dong, C. Ma, Q.J. Peng, E.-H. Han, W. Ke, Microstructure and stress corrosion cracking of a SA508-309L/308L-316L dissimilar metal weld joint in primary pressurized water reactor environment, J. Mater. Sci. Technol. (2019). https://doi.org/10.1016/j.jmst.2019.08.035

ur

na

[54] K. Chen, J.M. Wang, Z. Shen, D.H. Du, X.L. Guo, L.F. Zhang, P.L. Andresen, Effect of intergranular carbides on the cracking behavior of cold worked alloy 690 in subcritical and supercritical water, Corr. Sci. (2019) 108313. https://doi.org/10.1016/j.corsci.2019.108313.

Jo

[55] Z. Shen, K. Chen, D. Tweddle, G. He, K. Arioka, S. Lozano-Perez, Characterization of the crack initiation and propagation in Alloy 600 with a cold-worked surface, Corr. Sci. 152 (2019) 8292. https://doi.org/10.1016/j.corsci.2019.03.014 [56] S. Lozano-Perez, K. Kruska, I. Iyengar, T. Terachi, T. Yamada, The role of cold work and applied stress on surface oxidation of 304 stainless steel, Corr. Sci. 56 (2012) 78-85. https://doi.org/10.1016/j.corsci.2011.11.021 [57] B. Langelier, S.Y. Persaud, R.C. Newman, G.A. Botton, An atom probe tomography study of internal oxidation processes in Alloy 600, Acta Mater. 109 (2016) 55-68. https://doi.org/10.1016/j.actamat.2016.02.054 24

[58] T. Amadou, H. Sidhom, C. Braham, Double loop electrochemical potentiokinetic reactivation test optimization in checking of duplex stainless steel intergranular corrosion susceptibility, Metall. Mater. Trans. A 35 (2004) 3499-3513. https://doi.org/10.1007/s11661-004-0187-4 [59] G.S. Bai, S.P. Lu, D.Z. Li, Y.Y. Li, Intergranular corrosion behavior associated with delta-ferrite transformation of Ti-modified Super304H austenitic stainless steel, Corr. Sci. 90 (2015) 347-358.

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https://doi.org/10.1016/j.corsci.2014.10.031

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Figure captions Fig.1. Schematic drawing showing the mockup of the safe-end weld joint (a), how the bent beam specimens were extracted from the weld joint (b) and (c), and setup of the bent beam specimen and the fixture for the SCC test (d). In this figure, CPS represents colloidal silica slurry polishing, EPS represents electropolishing, MG represents mechanically grinding and M represent milling.

Fig. 2. OM and SEM observation of 308L weld cladding. (a), OM observation on cross-

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section of the cladding, (b), OM observation on surface of the cladding, and (c), SEM back-scattered electron image of γ/δ boundary.

Fig. 3. SEM observation of 308L weld cladding after the different surface treatments.

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(a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen,

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(c), milled specimen, and (d), electropolished specimen.

Fig. 4. Surface morphology of the 308L weld cladding after the different surface

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treatments analyzed 3D Profiler. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled specimen, and (d), electropolished

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specimen.

Fig. 5. Backscattered electron images of the cross-section of the 308L weld cladding

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after milling. (a), the overall microstructure, (b), the microstructure of the

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nanocrystalline layer, (c), backscattered electron image showing the elongated δ ferrite phase in the nanocrystalline layer, and (d), the microstructure of the underlying deformed layer.

Fig. 6. Load-depth curves of the δ ferrite and γ austenite phases on the 40 nm colloidal silica slurry polished, mechanically ground and milled 308L weld cladding obtained by nanoindentation tests. 26

Fig. 7. KAM distribution maps, the corresponding KAM value and microhardness distribution in the cross-section of 308L weld cladding after the different surface treatments. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled specimen, (d), electropolished specimen, (e), averaged KAM value vs. the distance from the surface, and (f), microhardness vs. the distance from the surface.

Fig. 8. Composition depth profiles of Cr, Ni and Fe on the surface of 308L weld

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cladding after the different surface treatments analyzed by XPS. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled specimen, and (d), electropolished specimen. (The Y-axis is the atomic percentage of

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Fe, Cr and Ni with different valence states in the sum of Fe, Cr and Ni elements)

Fig. 9. TEM observations and analyses of different regions of the milled 308L weld

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cladding. (a), TEM image of the phase boundary in bulk metal, (b), TEM image of the phase boundary in the underlying deformed layer, (c) and (d), TEM analyses of the

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nanocrystalline layer, (e), EDX mappings of the phase boundary, (f), composition profile across the carbide in the bulk metal, (g), composition profile across the γ/δ phase

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boundary in the bulk metal, and (h), composition profile across the elongated δ phase boundary in the nanocrystalline layer.

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Fig. 10. SEM observations of cracking on surface of the specimens with the different surface treatments following SCC tests in the simulated primary PWR environment at

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320 °C. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled specimen, and (d), electropolished specimen.

Fig. 11. SEM observations of crack initiation on cross-section of the specimens with the different surface treatments following SCC tests in the simulated primary PWR environment at 320 °C. (a), 40 nm colloidal silica slurry polished specimen, (b), 27

mechanically ground specimen, (c), (d) and (e), milled specimen, and (f), electropolished specimen.

Fig. 12. Summary of the concentration, maximum and average depth of the cracks on the cross-section of specimens with the different surface treatments following SCC tests in the simulated primary PWR environment at 320 °C.

Fig. 13. SEM observations and EBSD analysis of crack growth on cross-section of the

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milled specimen following SCC tests in the simulated primary PWR environment at 320 °C. (a), crack morphology on the side face of the specimen, (b), observation

showing crack growth along the γ/δ phase boundary but not along the γ/γ grain boundary (the location of this figure is marked by the while box in the insert figure ), (c),

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observation showing the difference preferential growth of crack along the γ/δ phase boundary intersected with deformation bands, (d), SEM observation of the crack at high

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magnification, and (e), EBSD analysis of a crack on cross-section of the specimen.

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Fig. 14. TEM observations and analyses of a crack that intersects with the elongated δ ferrite phase in the nanocrystalline layer. (a) and (b), crack tip morphology, (c),

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selected-area diffraction pattern of the oxide as marked by the red circle in figure (b), (d), composition profile across the oxidized δ ferrite phase, (e), composition profile along the oxidized γ/δ phase boundary, and (f), composition profile across the oxidized

ur

γ/δ phase boundary.

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Fig. 15. TEM observations and analyses of cracks that intersect with the carbide and austenite phase in the nanocrystalline layer. (a), crack tip morphology showing the crack intersected with the carbide, (b), crack tip morphology showing the crack intersected with the austenite phase, (c) and (d), composition profile along the oxidized δ ferrite phase as marked by the yellow lines in (a) and (b), respectively, and (e), EDX mappings of the crack tip area as marked by the red square in figure (b). 28

Fig. 16. Schematic drawing showing the role of δ ferrite phase in SCC of 308L weld cladding after machining in the simulated primary PWR environment. (a), surface microstructure of the 308L weld cladding, (b), formation of a uniform oxide film and preferential oxidation along the δ ferrite phase in the nanocrystalline layer, and (c), further oxidation and cracking along the phase boundary in the underlying deformed layer.

(a)

(b)

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lP

re

(c)

Machining marks

-p

ro of

(d)

Fig.1. Schematic drawing showing the mockup of the safe-end weld joint (a), how the bent beam specimens were extracted from the weld joint (b) and (c), and setup of the

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bent beam specimen and the fixture for the SCC test (d). In this figure, CPS represents colloidal silica slurry polishing, EPS represents electropolishing, MG represents

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mechanically grinding and M represent milling.

29

ro of -p

Fig. 2. OM and SEM observation of 308L weld cladding. (a), OM observation on cross-

re

section of the cladding, (b), OM observation on surface of the cladding, and (c), SEM

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na

lP

back-scattered electron image of γ/δ boundary.

30

ro of -p

Fig. 3. SEM observation of 308L weld cladding after the different surface treatments.

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(a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen,

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na

lP

(c), milled specimen, and (d), electropolished specimen.

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ro of -p

Fig. 4. Surface morphology of the 308L weld cladding after the different surface

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treatments analyzed 3D Profiler. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled specimen, and (d), electropolished

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na

lP

specimen.

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ro of -p

Fig. 5. Backscattered electron images of the cross-section of the 308L weld cladding

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after milling. (a), the overall microstructure, (b), the microstructure of the nanocrystalline layer, (c), backscattered electron image showing the elongated δ ferrite

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phase in the nanocrystalline layer, and (d), the microstructure of the underlying

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deformed layer.

33

Polished- Ground- Milled- Polished- Ground- Milled-

3000

Load (N)

2500 2000 1500 1000 500 0

0

25

50

75 100 125 150 175 200

Depth (nm)

ro of

Fig. 6. Load-depth curves of the δ ferrite and γ austenite phases on the 40 nm colloidal silica slurry polished, mechanically ground and milled 308L weld cladding obtained by

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na

lP

re

-p

nanoindentation tests.

34

ro of 250

Slurry polished Ground Milled Electropolished

re

0.8

(a) (b) (c) (d)

Hardness (HV)

Slurry polished Ground Milled Electropolished

0.6

0.4

200

lP

Average KAM value

-p

300

1.0

150

0.2

50-100

0-50

100-150 150-200 200-250

0

50

100

150

200

Distance from surface (m)

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Distance from surface (m)

Fig. 7. KAM distribution maps, the corresponding KAM value and microhardness distribution in the cross-section of 308L weld cladding after the different surface

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treatments. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled specimen, (d), electropolished specimen, (e), averaged

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KAM value vs. the distance from the surface, and (f), microhardness vs. the distance from the surface.

35

36

ro of

-p

re

lP

na

ur

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100

100 Cr hydroxide Cr oxide Cr (0) Ni (oxide) Ni (0) Fe (III) Fe (II) Fe (0)

75

50

25

20

40

60

80

75

50

25

0

100

20

100

Element concentration (at. %)

Cr hydroxide Cr oxide Cr (0) Ni (oxide) Ni (0) Fe (III) Fe (II) Fe (0)

75

50

25

40

60

80

Element concentration (at. %)

100

(c)

20

40

60

80

100

Sputtering time (s)

Sputtering time (s)

0

Cr hydroxide Cr oxide Cr (0) Ni (oxide) Ni (0) Fe (III) Fe (II) Fe (0)

Cr hydroxide Cr oxide Cr (0) Ni (oxide) Ni (0) Fe (III) Fe (II) Fe (0)

75

50

25

0

100

(d)

Sputtering time (s)

20

ro of

0

(b) Element concentration (at. %)

Element concentration (at. %)

(a)

40

60

80

100

-p

Sputtering time (s)

Fig. 8. Composition depth profiles of Cr, Ni and Fe on the surface of 308L weld

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cladding after the different surface treatments analyzed by XPS. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), milled

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specimen, and (d), electropolished specimen. (The Y-axis is the atomic percentage of

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Fe, Cr and Ni with different valence states in the sum of Fe, Cr and Ni elements)

37

(211)γ (110)γ (200)

(220)γ

(211) −

(011)

(f) Ferrite

60 40 20

Carbide

Austenite

-p

80

re

Atomic percentage (%)

100

ro of

[011]δ

Cr Fe Ni

0

Ferrite

80

Phase boundary

60

100

Austenite

40 20

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0 0

50

100 150 Distance (nm)

Atomic percentage (%)

(g)

na

Atomic percentage (%)

100

lP

0

Cr Fe Ni

200

50

100 150 Distance (nm)

200

250

(h) Ferrite

Austenite

80

Austenite

60

Cr Fe Ni

40 20 0

250

0

50

100

150 200 250 Distance (nm)

300

350

Fig. 9. TEM observations and analyses of different regions of the milled 308L weld

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cladding. (a), TEM image of the phase boundary in bulk metal, (b), TEM image of the phase boundary in the underlying deformed layer, (c) and (d), TEM analyses of the nanocrystalline layer, (e), EDX mappings of the phase boundary, (f), composition profile across the carbide in the bulk metal, (g), composition profile across the γ/δ phase boundary in the bulk metal, and (h), composition profile across the elongated δ phase boundary in the nanocrystalline layer. 38

39

ro of

-p

re

lP

na

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ro of

-p

Fig. 10. SEM observations of cracking on surface of the specimens with the different

surface treatments following SCC tests in the simulated primary PWR environment at

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320 °C. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground

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na

lP

specimen, (c), milled specimen, and (d), electropolished specimen.

40

ro of -p re lP na

Fig. 11. SEM observations of crack initiation on cross-section of the specimens with the different surface treatments following SCC tests in the simulated primary PWR

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environment at 320 °C. (a), 40 nm colloidal silica slurry polished specimen, (b), mechanically ground specimen, (c), (d) and (e), milled specimen, and (f),

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electropolished specimen.

41

120 Crack concerntration Maximum crack depth Average crack depth

Crack density/cm-1 Maximum crack depth/m Average crack depth/m

100 80 60 40 20 0

Milled

Electropolished

ro of

Slurry polished Ground

Fig. 12. Summary of the concentration, maximum and average depth of the cracks on

the cross-section of specimens with the different surface treatments following SCC tests

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ur

na

lP

re

-p

in the simulated primary PWR environment at 320 °C.

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Fig. 13. SEM observations and EBSD analysis of crack growth on cross-section of the milled specimen following SCC tests in the simulated primary PWR environment at

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320 °C. (a), crack morphology on the side face of the specimen, (b), observation showing crack growth along the γ/δ phase boundary but not along the γ/γ grain boundary

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(the location of this figure is marked by the while box in the insert figure ), (c), observation showing the difference preferential growth of crack along the γ/δ phase boundary intersected with deformation bands, (d), SEM observation of the crack at high magnification, and (e), EBSD analysis of a crack on cross-section of the specimen.

43

−

(111) (200) −

(111)

60

20

Ferrite

40

100

(f)

Austenite

150 200 250 Distance (nm) Oxide

60 40

300

20 0

350

Ferrite

O Cr Fe Ni

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O Cr Fe Ni

80

50

-p

60

0

Atomic percentage (%)

Atomic percentage (%)

80

O Cr Fe Ni

40

100 Oxide

Austenite

Oxide

0

−

(e)

Austenite

80

Maghemite/Magnetite[011] 100

(d)

ro of

Atomic percentage (%)

100

20

0

50

lP

0

100 150 200 Distance (nm)

250

300

0

50

100 150 Distance (nm)

200

250

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Fig. 14. TEM observations and analyses of a crack that intersects with the elongated δ ferrite phase in the nanocrystalline layer. (a) and (b), crack tip morphology, (c), selected-area diffraction pattern of the oxide as marked by the red circle in figure (b),

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(d), composition profile across the oxidized δ ferrite phase, (e), composition profile along the oxidized γ/δ phase boundary, and (f), composition profile across the oxidized

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γ/δ phase boundary.

44

100 Oxide

Carbide

(d)

Ferrite

80 60

O Cr Fe Ni

40 20 0

60

O Cr Fe Ni

40 20 0

50

100 150 200 Distance (nm)

250

0

50

100 150 200 Distance (nm)

250

re

-p

0

Austenite

Oxide

80

ro of

(c)

Atomic percentage (%)

Atomic percentage (%)

100

lP

Fig. 15. TEM observations and analyses of cracks that intersect with the carbide and austenite phase in the nanocrystalline layer. (a), crack tip morphology showing the

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crack intersected with the carbide, (b), crack tip morphology showing the crack intersected with the austenite phase, (c) and (d), composition profile along the oxidized δ ferrite phase as marked by the yellow lines in (a) and (b), respectively, and (e), EDX

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mappings of the crack tip area as marked by the red square in figure (b).

45

ro of -p

re

Fig. 16. Schematic drawing showing the role of δ ferrite phase in SCC of 308L weld cladding after machining in the simulated primary PWR environment. (a), surface

lP

microstructure of the 308L weld cladding, (b), formation of a uniform oxide film and preferential oxidation along the δ ferrite phase in the nanocrystalline layer, and (c), further oxidation and cracking along the phase boundary in the underlying deformed

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layer.

Table 1 Chemical composition (wt.%) of 308L weld metal Si

Mn

P

S

Cr

Ni

Cu

Nb

Fe

0.016 0.015

0.32 0.18

1.33 1.53

0.015 0.014

0.011 0.011

19.88 19.29

10.30 9.92

0.065 0.05

0.01 0.01

Bal. Bal.

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308L welding wire 308L cladding

C

Table 2 Surface microhardness and roughness of the specimens after different surface treatments Surface treatment

Microhardness, Hv

Roughness (Ra), μm

CPS

209.2±11.7

0.48

EPS

206.4±10.9

0.96

MG

376.9±27.0

6.66 46

M(peak)

412.1±16.7

M(valley)

358.8±20.1

3.29

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na

lP

re

-p

ro of

CPS: colloidal silica slurry polishing, EPS: electropolishing, MG: mechanically grinding, M: milling.

47