Role of residual ferrites on crevice SCC of austenitic stainless steels in PWR water with high-dissolved oxygen

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Original Article

Role of residual ferrites on crevice SCC of austenitic stainless steels in

PWR water with high-dissolved oxygen

Abdullah Sinjlawi

a

, Junjie Chen

a

, Ho-Sub Kim

a,c

, Hyeon Bae Lee

a

, Changheui Jang

a,*

,

Sanghoon Lee

d

a_{Korea Advanced Institute of Science and Technology, Daejeon, 34141, Republic of Korea} c_{Defense Agency for Technology and Quality, Jinju, Gyeongsangnam-do, 52851, Republic of Korea} d_{Korea Institute of Materials Science, Changwon, Gyeongsangnam-do, 51508, Republic of Korea}

a r t i c l e i n f o

Article history:

Received 4 February 2020 Received in revised form 15 April 2020

Accepted 23 April 2020 Available online 28 April 2020

Keywords: Residuald-ferrite Austenitic stainless steel Crevice beam bent (CBB) Stress corrosion cracking (SCC)

a b s t r a c t

The crevice stress corrosion cracking (SCC) susceptibility of austenitic stainless steels was evaluated in simulated pressurized water reactor (PWR) environments. To simulate the abnormal condition in tem-porary clamping devices on leaking small bore pipes, crevice bent beam (CBB) tests were performed in the oxygenated as well as hydrogenated conditions. No SCC cracks were found for SS316 in both con-ditions. SS304 also showed good resistance in the hydrogenated condition. However, all SS304 specimens showed SCC cracks in the oxygenated condition, indicating poor crevice SCC resistance. It was found that residual ferrites were selectively dissolved because of the galvanic corrosion coupled with the neigh-bouring austenite phase, resulting in SCC initiation in SS304. Crack morphologies were mostly trans-granular assisted by the damagedd-ferrite and deformation-induced slip bands.

1. Introduction

There are many small-bore pipes made of austenitic stainless steels (ASS) in the primary coolant systems of the pressurized water reactors (PWRs). During the reactor operation, these pipe struc-tures including valves and joints are subjected to various vibra-tional and cyclic loads, which may cause the failure of the welded joints due to high cycle fatigue. Such high cycle fatigue failures of small-bore pipe components have been widely reported over de-cades [1e4], raising the safety concern of the operating nuclear power plants [2]. When failure or leakage is found in these small components, temporary maintenance strategy is needed to avoid the unscheduled shutdown [5]. As a temporary measure, a clamping device could be installed surrounding the damaged part to contain the leaked primary water and prevent the further leak of primary water, thus allowing the normal operation until the next scheduled shutdown. The clamping device should be made of ASSs to match the damaged small-bore pipe structures.

Meanwhile, it is anticipated that the inside of clamping device will be eventually ﬁlled with the leaked primary water during

normal operation, resulting in a stagnant condition. Also, because of the internal pressure, local tensile stress and crevices would be developed in the bolted regions. An additional concern is that, during the clamping process, air will be trapped and dissolved in the leaked coolant, creating an oxygenated water environment within the clamping devices. In the PWRs, higher content of dis-solved oxygen (DO) remained in the occluded areas found to in-crease the electrochemical potential (ECP) to higher levels [6,7] and the inward oxidant diffusion towards the crevice [8,9]. Also, the stress corrosion cracking (SCC) susceptibility of the ASSs would be increased in the high DO condition [6,7,10]. Thus, there is concern on the integrity of the materials of clamping devices in crevice environment of high DO PWR water, especially the SCC suscepti-bility of ASSs. Therefore, in this study, the crevice SCC behavior of ASSs was investigated in oxygenated (8 ppm DO) and hydrogenated PWR primary water environments. The crevice condition was simulated using crevice bent-beam (CBB) specimens. The SCC cracks were analyzed using various electron microscope tech-niques. Finally, the effect of materials and microstructure on SCC initiation and propagation was discussed.

* Corresponding author.

E-mail address:[email protected](C. Jang).

Contents lists available atScienceDirect

Nuclear Engineering and Technology

j o u r n a l h o me p a g e : w w w .e l se v i e r. co m/ lo ca t e / n e t

https://doi.org/10.1016/j.net.2020.04.023

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2. Materials and method 2.1. Materials

Two commercial ASSs, SS304 and SS316, were used in this study. Both materials were solution annealed at 1000C for 30e60 min, followed by water quenching. The chemical compositions were analyzed by the inductively coupled plasma (ICP) technique, and the results are summarized inTable 1. It should be noted that small amount of non-uniformly dispersed residual ferrites is present in the commercial alloys used in this study (9% for SS304 and 2% for SS316 according to Schaefﬂer diagram). Also, the tensile properties of the materials were measured at room temperature and the re-sults are summarized inTable 2.

2.2. Preparation of CBB specimens and test method

The evaluation of the SCC initiation resistance was performed by a constant load method using a CBB design. The dimensions of specimens and CBB_{fixture (made of 304SS) are shown in}Fig. 1. CBB test method has been widely used because of its ability to create high crack density with large sizes [11e14]. Moreover, it can closely simulate the desired crevice condition through graphite_{fiber wool} which was placed between the upper jig and specimen and it serves as crevice creator. All specimens were machined from large rectan-gular blocks, and surfaces were mechanically ground using SiC pa-pers down to 4000 grits. Afterwards, specimens and CBBfixtures were rinsed and cleaned by demineralised water and ethanol using an ultrasonicator before being assembled. Bending loadings equiv-alent to 2e3% strain was applied on the specimens using a

screw-type Instron 4204. From the stress strain curves, the applied stress was 352 MPa for SS304 and 318 MPa for SS316, respectively. Sub-sequently, all parts were assembled (jigs, specimens, and the graphite wool) by tightening bolts at both sides. A torque wrench has been used at the lowest torque value (30 kgf-cm) in order to ensure a suitable gap between the jigs. A Vernier caliper was used to measure the gap at four different points at each face of the lower and upper jigs. Considering the specimen thickness (2 mm) and graphite wool thickness (0.1_{e0.2 mm), the gap between jigs was kept at} 2.3e2.5 mm, such that water can enters and remains in stagnant condition between specimen surface and graphite wool. It should be noted that the exact local environment between specimen surface and graphite wool could not be identiﬁed due to space limitation. Two CBB ﬁxtures were used with different jig curvature (50 and 100 mm radii) to analyze the effect of applied stress.

2.3. CBB test conditions

The test temperature and pressure were at 325C and 15.5 MPa, respectively, simulating the primary coolant system water in the PWRs. The pH was kept at 6e7 by recirculating loop system. For oxygenated water condition, artiﬁcial air mixture of pure oxygen (20 vol%, 99.99% purity) and pure nitrogen (80 vol%, 99.99% purity) was continuously bubbled into the testing loop to achieve the target DO content (~8 ppm). For hydrogenated water condition, pure hydrogen (99.99%) was continuously injected into the loop system and with a slight overpressure in water tank to maintain the target DH content (2.65 ppm). The DO and DH contents, water conductivity, and pH value were continuously monitored during the entire test period. The specimens have been exposed for 500 h in hydrogenated PWR water condition and 1000 h in oxygenated PWR water condition, respectively. The test matrix showing ma-terials and test conditions are summarized inTable 3.

2.4. Microstructure analysis

After the test, top surface of specimens was observed under scanning electron microscopy (SEM) to detect cracks. Specimens

Table 1

The chemical compositions of the austenitic stainless steels used in this study (in wt.%).

Material Fe Cr Ni Si Mn C P Co S Mo Ferrite content

SS304 Bal. 18.10 8.05 0.40 0.85 0.057 0.032 0.23 0.0036 e ~9% SS316 Bal. 16.80 11.00 0.57 0.93 0.07 0.015 e 0.004 2.58 ~2%

Table 2

The tensile properties of the austenitic stainless steels used in this study. Material Ultimate Tensile Strength, MPa Yield Strength, MPa Elongation,

%

SS304 729 304 82

SS316 601 265 83

Fig. 1. Dimensions of CBBﬁxture and specimen used in this study.

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with long cracks were sectioned into small pieces using a low-speed saw to minimize damage to any cracks. The cross-sections were ground starting from 1200 grit SiC paper for several mi-nutes until the evidence of crack starts to appear and observed under an optical microscope (OM). Then the area was carefully

ground using 4000 grit SiC paper until the whole crack shape was revealed. The detailed analysis of cracks and surrounding micro-structure, transmission electron microscopy (TEM), focused ion beam (FIB), and electron backscattering diffraction (EBSD) were used. Meanwhile, to investigate surface condition of the tested

Table 3

CBB test materials and test conditions.

Condition Specimen Exposure time (h) Number of specimens Water Chemistry Hydrogenated SS304 SS316 500 2 DO< 5 ppb, DH ~ 2.65 ppm 2 Oxygenated SS304-DO 1000 6 DO ~ 8 ppm, DH< 5 ppb SS316-DO 6 Table 4

CBB test results of SS304 and SS316 immersed in hydrogenated and oxygenated PWR water conditions.

Hydrogenated PWR (<5 ppb DO) Oxygenated PWR (~8 ppm DO)

Material CBB design No. of cracksa _{Max. length} _Material _{CBB design} _{No. of cracks/specimen*} _{Max. length}

SS304 50-R 0 SS304 50-R 6 1800mm

100-R 0 e 100-R 7 1300mm

SS316 50-R 0 e SS316 50-R 0 0

100-R 0 e 100-R 0 0

Average crack length ~400mm.

a_{Only cracks with length}_>100_m_{m have been considered as SCC.}

Fig. 2. TEM result of the oxygenated water condition for (a) SS316 and (b) SS304. TEM-EDS line scan showing atomic percentage of surface oxides elements of (c) SS316 and (d) SS304.

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specimens, surface oxides were removed by light manual-grinding using 4000# SiC grit paper. In summary, the following analytical tools have been used in this study:

1. Hitachi SU-5000 SEM with energy dispersive spectroscope (EDS) and back-scattered electron (BSE) detectors for crack scanning.

2. Helios G4 Ultra High-Resolution focused ion beam (FIB) with Schottky thermal ﬁeld emitter electron source and Gallium liquid metal ion source for TEM samples preparation.

3. Talos F200X TEM at 200 kV accelerating voltage with EDS de-tector for oxide and crack tip analysis.

4. X'Pert-PRO MRD with the maximum power of 3 kW High-Resolution thin-ﬁlm X-Ray (XRD).

3. Results and discussion 3.1. CBB test results

The CBB test results with crack quantiﬁcation are summarized inTable 4. It is clear that both SS304 and SS316 showed good SCC resistance in the hydrogenated PWR water condition such that cracks were not found in any specimens. However, in the oxygenated PWR water condition, SCC resistance is quite different for two ASSs. While, SS316 was still resistant to SCC (showing no cracks), SS304 showed high susceptibility to the crevice SCC in the oxygenated PWR water condition such that all

SS304 specimen showed multiple cracks as large as 1800

m

m. In this study, cracks longer than 100

m

m was considered SCC cracks and counted. There was no significant quantitative difference in SCC crack was noticed between the two CBB design (50-R and 100-R). The small difference in the applied strain (~1%) could be considered not sufficient to cause significant difference in the tensile stress pro_{file at the outer surface of specimens that was in} the plastically deformed condition. It should be noted that, as cracks were only found in the oxygenated PWR water condition, specimens exposed to such a condition were further analyzed in the following sections.

3.2. Surface oxide analysis

The results of TEM/EDS analysis of SS316 and SS304 exposed in the oxygenated PWR water condition are shown inFig. 2. According to EDS line scan results (Fig. 2c and d) and selected area diffraction (SAD) patterns (Fig. 2a and b), a duplex oxide structure consisted of the outer Fe-rich hematite and inner Cr-rich spinel oxide inner layers for both materials. For SS316, small amount of Ni was present in both oxide layers, suggesting the inner oxide layer would be nickel ferrite spinel, which is known to have higher stability over other spinel structures in the oxygenated water condition [15,16]. On the other hand, Ni content was mostly negligible in the oxide layer formed on SS304.

Moreover, at the interface between the oxide layer and matrix, dark spots were observed for SS304 (Fig. 2b). EDS result showed

Fig. 3. (a) Surface of SS304 exposed to the oxygenated PWR water condition showing damaged oxide layers. Detailed view of damaged areas after oxide removal showing pit-like structure in (b) austenite matrix and (c)d-ferrite surrounded by austenite.

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that these spots are highly enriched with Ni but deprived of Cr and Fe (Fig. 2d). It is interesting that oxygen content in these spots is very low, but oxygen is present below the Ni-rich layer. Further-more, it was found that hematite co-existed with spinel in the inner oxide layer for SS304 (Fig. 2b), which could be attributed to the easier oxygen penetration through less protective oxide layer of

SS304. The difference in the protectiveness of oxide layers of both materials could be consistent with the observed average oxide thickness (~700 nm for SS304 vs 350 nm for SS316) in the oxygenated PWR water condition. The observed oxide character-istics in the high DO environment are in good agreement with the previously published reports [3e5,9,17,18].

Fig. 4. (a) Surface of SS304 exposed to the oxygenated PWR water showing selective dissolution ofd-ferrites and associated microcracks, (b) the magniﬁed view ofd-ferrite with microcracks within, (c) EDS analysis results for points 1 and 2 in (b), and (d) cracks initiation at the interface of austenite matrix andd-ferrite (org/dinterphase boundary).

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The surface of SS304 specimen tested in the oxygenated PWR water condition was shown inFig. 3. There are many areas showing damaged oxide layer in both austenite matrix and residual

d

-fer-rites. Though not shown here, damaged oxide layer was not observed for SS316 specimens tested in the oxygenated PWR water condition. The localized attack could be associated with the oxide ﬁlm defects [18,19], which could act as active channels of oxygen diffusion [20,21] and eventually cause oxideﬁlm breakdown. The detailed observation of the damaged area shows the pit-like for-mation (Fig. 3c and d taken after surface oxide removal) due to the localized attack. Though pit-like structures were present in both austenite matrix (Fig. 3b) and

d

-ferrites (Fig. 3c), the degree of attack is clearly more signiﬁcant for

d

-ferrite surrounded by austenite matrix. The selective anodic dissolution of more active ferrites galvanically coupled with austenite matrix could have caused much greater pit-like defects in ferrites. It should be noted that there are micro-cracks at the bottom of pit-like structure in

d

-ferrites. The detailed analysis of cracks is described in the following

section. Meanwhile, the graphiteﬁber wool role in the SCC process is a matter of discussion. It was suggested that bicarbonate ions (HCO3 ) were formed inside the crevice of CBB specimen [12], which would raise the local water conductivity and eventually accelerate the SCC initiation [22]. Moreover, the dissolved metal species could have been increased due to the high ferric ion products, the strong oxidizer. This increase in the ferric ion prod-ucts might altered the electrochemical activity of the solution [23]. Nonetheless, the local water chemistry condition in this study cannot be compared with the BWR environment where dissolved metal species level is low to ensure a low corrosion potential [16]. 3.3. Crack analysis

3.3.1. The role of microstructure on the SCC initiation and propagation

The residual

d

-ferrite usually remained in the form of globules or continuous islands in austenitic stainless steels. In this study,

Fig. 5. Examples ofd-ferrite role on the SCC crack propagation in SS304 exposed to the oxygenated PWR water; (a) surface crack associated withd-ferrites (before oxide removal), (b) surface crack associated withd-ferrites (after oxide removal), (c) crack branching neard-ferrite along main crack path, and (d) EDS analysis of points in (c).

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these

d

-ferrites have a crucial role in the SCC initiation and prop-agation as can be seen inFigs. 4e6. As shown inFig. 4a-d, the

d

-ferrite globules are severely damaged and several mircocracks have been initiated at the interface of austenite matrix and

d

-ferrite (or

d

/

g

interphase boundary) and propagated towards either the

d

-ferrite phase (Fig. 4a and b) or the austenite matrix (Fig. 4d). Meanwhile, large

d

-ferrite globules are found along the long surface cracks as shown in Fig. 5a (before oxide removal) and 5b (after oxide removal). As

d

-ferrites were already damaged with micro-cracks, they could be the preferential path for the crack

propagation, resulting in the representative crack morphologies shown inFig. 5a and b. In addition,

d

-ferrites provided sites of crack branching from the main crack as shown inFig. 5c. Thus, both the propagation of the main crack and crack branching were associated with

d

-ferrites for CBB tested SS304 in oxygenated PWR water condition. The cross-sectional view of the main SCC crack is shown inFig. 6a. As shown in the ﬁgure, though the crack propagated mostly through austenite matrix, there are more

d

-ferrite globules along the crack path compared to surrounding region. This is consistent with the observation on the surface crack inFig. 5. From

Fig. 6. (a) A cross-sectional view of the crack inFig. 5a and (b) EBSD-IPF image of the rectangular box in (a) showing the TGSCC dominated morphology, and (c) EBSD phase map showing ferrites (in green). (For interpretation of the references to colour in thisﬁgure legend, the reader is referred to the Web version of this article.)

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the EBSD image (Fig. 6b), it is clear that crack propagated pre-dominantly in TG mode but some TG to IG transition was observed along the crack propagation path. In EBSD phase map (Fig. 6c), the residues of

d

-ferrites could be observed on the wall of the crack, which indicate the possible involvement of

d

-ferrites in the crack

propagation. Based on the crack morphology observation on the surface and cross-section of CBB tested SS304 in oxygenated PWR water condition, it could be said that SCC cracks of SS304 mostly propagated in TGSCC cracking mode aided by the presence of the

d

-ferrite globules along the path.

Fig. 7. (a) TEM dark-ﬁeld image of the damagedd-ferrites along the main SCC crack, (b) EDS elemental mapping results, (c) High magniﬁcation images of area marked in (a) showing microcracks branched into adjacent austenite grains, and (d) and (e) EDS elemental scan results of marked lines in (a).

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Fig. 8. (a) TEM dark-ﬁeld tem image of one of the microcracks branched into austenite grain (fromFig. 7c) and EDS line scan results of marked lines (b) A-B, (b) C-D, and (c) E-F, respectively.

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The role of

d

-ferrites on the TGSCC cracking mode could be explained as follows. As mentioned in the previous section, a micro-galvanic cell could be formed between the austenite and

d

-ferrite (particularly at

d

/

g

interphase). Such a galvanic corrosion led to a signiﬁcant selective (anodic) dissolution in the

d

-ferrite region because of the large area portion difference between the two phases. As a result, the elements dissolution of

d

-ferrite (primarily Fe) would signiﬁcantly deteriorate the phase stability and increase the SCC susceptibility in the oxygenated water [24]. In addition, the microcracks in

d

-ferrite (Fig. 4b) provide the oxygen diffusion path towards the austenite grain, resulting in localized attack in the adjacent austenite grains. To conﬁrm such mechanism, a TEM analysis was performed at different locations that cover

d

-ferrite,

d

/

g

interphase, and austenite grains. Fig. 7a shows the

d

-ferrites located along the main SCC crack and microcracks branched into adjacent austenite grain. The elemental mapping of the area is shown inFig. 7b. EDS line-scan results (Fig. 7d and e) conﬁrm that the

d

-ferrites were selectively oxidized while both Fe and Ni was practically dissolved from the

d

-ferrite, leaving Cr-rich oxides. In the magniﬁed images of the marked areas inFig. 7a, many micro-cracks were formed in the damaged

d

-ferrites and the adjacent austenite grains (Fig. 7c). In addition, EDS line-scan was performed on one of the microcracks branched into adjacent austenite grain (second image inFig. 7c) and the results are shown inFig. 8. It is clear that both Fe and Ni were selectively dissolved while Cr was enriched at the tip and oxidized, which is consistent with the previous studies [9,14,25e27].

To determine the structures of Cr-rich oxides formed at the cracks, three different locations (A, B, and C) have been selected in Fig. 7a and diffraction patterns were obtained. As shown inFig. 9, amorphous oxides (CreO) were formed at the oxidized

d

-ferrite globule (points A and C) where Fe was mostly dissolved (Black spots

are remaining Fe). Meanwhile, a possible nano-crystalline Cr2O3

structure was observed at point B which is the surface oxide layer and considered to be exposed to high DO water. It is noteworthy that the Cr-rich amorphous oxides were formed due to the selective dissolution of Fe and Ni from the

d

-ferrites galvanically coupled with austenite. The amorphous oxide (CreO) is known to be highly susceptible to rupture [19] due to its brittle nature. As a result, many microcracks were formed in the amorphous CreO (oxidized

d

-ferrite globule) and at the interface with austenite grains (Fig. 7a). 3.3.2. The role of slip bands

Slip bands near the SCC cracks are shown inFig. 10. It has been reported that the slip bands usually provide a preferential and fast oxygen diffusion path, but the mechanism is complicated as it is related to the stress concentrators and dislocations motion in addition to their alignment with the crack propagation plane [21,25,28e30]. In this study, both oxidation and crack branching along slip bands were observed (Fig. 10a), suggesting the slip bands provide easier diffusion path of oxygen, preferential oxidation, and crack propagation path. On the other hand, slip bands are respon-sible for oxideﬁlm breakdown because of the local increase of the dislocation density at these areas. It is noteworthy that the large amount of slip bands and twins observed in the plastically deformed SS304 (CBB specimen used in this study) are attributed to its low stacking fault energy [21,31]. Moreover, the higher density of slip bands is also attributed to the microstructural difference be-tween austenite and

d

-ferrite such that the stress-strain distribu-tion is non-uniform along the

d

/

g

interphase boundary [23] which could concentrate the slip bands around these areas as shown in Fig. 10b. The higher density of slip bands around the

d

-ferrites and the presence of microcracks in the damaged areas around them would further accelerate the oxidation of slip bands and

Fig. 10. (a) Oxidation and crack branching along the slip bands and (b) oxidized slip bands in the austenite matrix and neard-ferrite globules.

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transgranular crack propagation, which is consistent with the predominantly TGSCC cracking mode described in the previous section.

3.3.3. SCC crack growth in the austenite phase

To compare the SCC crack growth mechanism in austenite grain (without the effect of

d

-ferrite) with that associated with the damaged

d

-ferrites, one of crack tips in the austenite region was machined and analyzed.Fig. 11a shows the machine crack tip in the austenite phase and the area of EDS elemental mapping analysis. EDS elemental mapping results (Fig. 11b) shows that the surfaces of cracks are covered with Cr-rich oxides. EDS line scan results (Fig. 11c and d) show that both Fe and Ni were selectively dissolved while Cr was enriched at the tip and oxidized, which is consistent with the previous observation (Fig. 8). Moreover, TEM analysis shows the presence of Cr-rich amorphous oxides at the crack tip and crystalline oxides behind the tip (Fig. 12a and b), which is consistent with the oxide structure formed near the damaged

d

-ferrites (Fig. 9). Mean-while, several straight sections are found along the crack path, indicating the possible role of slip bands. The evidence of crack propagation by slip-rupture and oxidation [24] could be seen in the HRTEM image of the crack tip (Fig. 12c). The presence of the narrow section between the main crack tip and advancing crack could indicate the recently ruptured oxide layer. Furthermore, the near straight shape of the advancing crack could implicate the involve-ment of slip band in the crack propagation in austenite grain.

4. Conclusions

The crevice SCC susceptibility of SS304 and SS316 was evaluated in simulated pressurized water reactor (PWR) environments with high oxygen content. To simulate the crevice conditions, crevice bent-beam (CBB) specimens were used for SCC tests in the oxygenated as well as hydrogenated PWR water conditions. Based on the test results and subsequent analysis, the following conclu-sions were drawn.

1. No SCC cracks were found for SS316 in both conditions. SS304 also showed good SCC resistance in the hydrogenated condition. However, all SS304 specimens showed SCC cracks in the oxygenated condition, indicating poor crevice SCC resistance. 2. In SS304 tested in the oxygenated PWR condition, residual

fer-rites were selectively dissolved because of the galvanic corro-sion coupled with the neighbouring austenite phase, resulting in SCC initiation in ferrites and adjacent austenite grains. 3. Crack propagation was mostly transgranular cracking assisted

by the damaged

d

-ferrite and slip bands along the path for SS304 tested in the oxygenated PWR condition.

4. The higher density of slip bands around the

d

-ferrites and the presence of microcracks in the damaged areas around them further accelerated the oxidation of slip bands and transgranular crack propagation.

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Data availability

The raw/processed data required to reproduce these ﬁndings cannot be shared at this time due to technical or time limitations. They will be shared upon request.

Declaration of competing interest

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

appeared to inﬂuence the work reported in this paper.

Acknowledgement

This study was supported by the Energy R&D Program (No. 20161510200340) of the MOTIE/KETEP of the Republic of Korea, the Engineering Research Center Program (No. 2016R1A5A1013919) and the Nuclear R&D Program (No. 2019M2D2A2050927) of MSIT/NRF of the Republic of Korea. Financial support for three of the authors was provided by the BK-Plus Program of the MSIT of the Republic of Korea.

Fig. 12. TEM analysis of crack tip in austenite grain: (a) TEM image of crack tip, (b) high magnification image and FFT patterns of point 1 in (a), and (c) high magnification image of point 2 showing SCC advance with slip rupture mechanism and high magnification image and FFT patterns of point 2 in (a).

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