Characterization of microstructure and local deformation in 316NG weld heat-affected zone and stress corrosion cracking in high temperature water

Corrosion Science 53 (2011) 1916–1932

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Characterization of microstructure and local deformation in 316NG weld heat-affected zone and stress corrosion cracking in high temperature water Zhanpeng Lu, Tetsuo Shoji ⇑, Fanjiang Meng 1, He Xue, Yubing Qiu, Yoichi Takeda, Koji Negishi Fracture and Reliability Research Institute, Faculty of Engineering, Tohoku University, Aramaki Aoba 6-6-10, Aoba-ku, Sendai 980-8579, Japan

a r t i c l e

i n f o

Article history: Received 18 November 2010 Accepted 6 February 2011 Available online 12 February 2011 Keywords: A. Stainless steel B. SEM Stress corrosion C. Welding C. Reactor condition

a b s t r a c t Microstructure and local deformation in 316NG weld heat-affected zones were measured by electronback scattering diffraction and hardness measurements. With increasing the distance from the fusion line, kernel average misorientation decreases and the fraction of R3 boundaries increases. Stress corrosion cracking growth rates in high temperature water were measured at different locations in the heat-affected zones that correspond to different levels of strain-hardening represented by kernel average misorientation and hardness distribution. Intergranular cracking along random boundaries as well as extensive intergranular crack branching is observed in the heat-affected zone near the weld fusion line. Ó 2011 Elsevier Ltd. All rights reserved.

1. Introduction In recent years, stress corrosion cracking (SCC) has been experienced in boiling water reactor (BWR) components fabricated with low-carbon stainless steels (SS), where cracking was observed in primary loop re-circulation pipes (PLR) and core shrouds [1–6]. Cracking in PLR pipes was frequently observed near the welds inside the piping, where transgranular cracking was initiated from the inner surface and propagated inside the pipe material by intergranular stress corrosion cracking. Cracking in PLR pipes are mainly confined in the heat-affected zone (HAZ) while there are some cases that cracks penetrated into the weld metal [1–4]. It was also found that there were cases that cracks penetrated into the weld metal in core shroud welds. The cracking behavior in the area near the fusion line is of importance for the safety evaluation of PLR pipes. Itow and Suzuki [7] have investigated the SCC growth behavior in the heat-affected zone near the fusion line in a 316NG (base)–316L (weld metal) weld. Their results show that the angle of the notch direction against the weld metal dendrite direction, and the orientation of the microstructure against the applied stress affect the crack growth or retardation behavior near the fusion boundary. Extensive crack branching in the HAZ near the fusion line is observed. Arai et al. [8] also show that there is extensive stress corrosion crack branching in the HAZ near the fusion boundary in a 316L (base)–316L (weld metal) weld, which ⇑ Corresponding author. Tel.: +81 22 7957517. E-mail address: [email protected] (T. Shoji). On leave from Institute of Metal Research, Chinese Academy of Science, Shenyang, China. 1

0010-938X/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2011.02.009

would inhibit the crack penetration into the weld metal. The distribution of the residual stress in the HAZ would have a significant effect on the cracking path. Extensive analyses on these cracked components fabricated from non-sensitized SS show that residual stress and strain are crucial for the high crack growth rates observed in BWR plants [1–6,9–13]. Angeliu et al. [5] used transmission electron microscope to assess the dislocation density and carbides, and used electron-back scattered diffraction (EBSD) to quantify the residual strain by measuring the intra-grain misorientation in the 304L SS weld HAZ. Jiao and Was [14] reported that irradiation assisted stress corrosion cracking of austenitic stainless steel initiated at locations where a large slip channel intersects a grain boundary in simulated BWR environment, showing that local deformation may play a key role in the irradiation assisted stress corrosion cracking in light water reactor core component. Alexandreanu et al. [15] have investigated the effect of grain boundary character distribution on the high temperature creep of Ni–16Cr–9Fe alloys at 360 °C. Their results show that extrinsic grain boundary dislocations (EGBDs) adsorption are annihilated at high angle boundaries (HABs) at a rate that is on average three times that at coincident site lattice boundaries (CSLBs), implying that a grain boundary diffusion coefficient in CSLBs that is 12 times lower than that in HABs. They also found that the hardness in the vicinity of CSLBs is higher than that in HABs, and the grain averaged hardness increases with the fraction of contiguous CSLBs. These results show that CSLBs impede dislocation adsorption into the grain boundary, thereby increasing lattice hardening and internal stress in the sample, resulting in a reduced creep rate. Alexandreanu and Was [16] also reported that 61% of the boundaries that deformed in argon at 360 °C also

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Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932 Table 1 Chemical composition (wt%) of 316NG SS base metal used in fabrication of pipe welds. C

Si

Mn

P

S

Ni

Cr

Mo

N

Fe

0.01

0.48

1.42

0.023

0.003

11.64

17.50

2.08

0.116

Bal.

cracked in 360 °C simulated pressurized water reactor (PWR) primary water, while only 17% of cracked boundaries showed no prior deformation in argon. These results correlated grain boundary deformation and IGSCC of Ni-base alloys in high temperature water. Lozano-Perez et al. [17] have conducted multi-scale study of SCC in cold-worked 304 SS in high temperature water. The oxidation of deformation bands and related cracking paths in coldworked stainless steel were characterized and discussed. Shoji [18] has proposed two fundamental driving forces for stress corrosion cracking in high temperature water: oxidation localization and oxidation acceleration in which local deformation would play a crucial role. The effects of grain boundary characteristic on SCC of stainless steels in simulated BWR environments have been reported [19–21]. The effects of water chemistry on SCC growth of 316L HAZ and weld metals have been investigated [12,22]. In the

present work, the microstructure and local deformation in 316L weld HAZ material from a pipe weld mockup are determined by electron-back scattering diffraction (EBSD) technique and hardness measurements. Stress corrosion cracking growth rate in HAZs with various distances from the fusion line in simulated BWR water environments are measured. The roles of local deformation and microstructure in SCC are analyzed. 2. Material and experimental methods 2.1. Materials and specimens preparation Type 316NG pipe (600 mm diameter, 41.27 mm wall thickness) was welded using two heats of Type 316L weld wire [23]. The chemical compositions of 316NG PLR pipe are shown in Table 1. The high ferrite content wire composition in wt% was 0.015 C, 1.75 Mn, 0.35 Si, 0.014 S, 0.017 P, 19.2 Cr, 12.3 Ni, 2.61 Mo, and 0.05 Cu, which yields a Cr(eq) of 22.33% and a Ni(eq) of 13.62%. The low ferrite content wire content composition was 0.022 C, 1.85 Mn, 0.44 Si, 0.001 S, 0.021 P, 19.34 Cr, 12.68 Ni. 2.51 Mo, and 0.26 Cu, which yields a Cr(eq) of 22.51% and a Ni(eq) of 14.26%. The nominal ferrite content for the two welds was 10% and 8%, respectively. The ferrite compositions correspond to ferrite

Fig. 1. Observed locations of crack tips on the side surfaces of specimens WOH1 and HTH2 after the phase-1 SCC test. Side grooves were removed and the side surfaces were polished before observation. (a) Specimen WOH1 and (b) specimen HTH2.

Table 2 Testing procedures for two 316NG SS HAZ specimens in simulated BWR environments.

#

Phase

Step No.

DO (ppm)

Specimens size, stress intensity factor K# and loading mode

Test period (cycles or h)

Phase-1

F11 F12 F13 C11 C12

2 2 2 2 0.2

288 cycles 576 cycles 864 cycles 1157 h 150 h – –

Phase-2

F21 F22 F23 C21

2 2 2 2

Tri. loading, 0.01 Hz, R = 0.3, Kmax (WOH1)  34 MPa m0.5, Kmax (HTH2)  36 MPa m0.5 20 mm thick, Tri. loading, 0.01 Hz, R = 0.5, Kmax  34 MPa m0.5, Kmax (HTH2)  36 MPa m0.5 Tri. loading, 0.01 Hz, R = 0.7, Kmax  34 MPa m0.5, Kmax (HTH2)  36 MPa m0.5 20 mm thick, constant loading, K  34 MPa m0.5, K (HTH2)  36 MPa m0.5 Thick, constant loading, K (WOH1)  34 MPa m0.5, K(HTH2)  36 MPa m0.5 Test interrupted. The side grooves were removed. One piece of slice in each specimen was taken out Pre-cracking in air (second pre-crack) to make the crack tip near the fusion line, R = 0.2, Kmax (WOH1)  17 MPa m0.5, Kmax (HTH2)  19 MPa m0.5 Tri. loading, 0.01 Hz, R = 0.3, Kmax (WOH1)  34 MPa m0.5, Kmax (HTH2)  41 MPa m0.5 Tri. loading, 0.01 Hz, R = 0.5, Kmax (WOH1)  34 MPa m0.5, Kmax (HTH2)  41 MPa m0.5 Tri. loading, 0.01 Hz, R = 0.7, Kmax (WOH1)  34 MPa m0.5, Kmax (HTH2)  41 MPa m0.5 Constant loading, K  41 MPa m0.5

Values of K at different locations of the specimens are calculated with finite element methods.

288 cycles 576 cycles 864 cycles 1434 h

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Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932 Block A for fracture surface

Specimen thickness

1

Phase-1 SCC test

18 1.0

men WOH1 was prepared from a weld prepared with a low ferrite content wire. The design of the specimen has been described in previous publications [10,12]. The notch positions in two specimens were different for testing SCC growth behavior at different locations in the heat-affected zones, which are shown in a part of SCC tests. The distribution of hardness in the tested specimens was measured on the SCC tested specimens with Shimazu Vickers hardness tester with a loading of 2.45 N and a holding time of 15 s. More complete distribution of Vickers hardness (HV1) in the HAZ was measured with a loading of 9.8 N and a holding time of 15 s on an as-welded HAZ block prepared by a high ferrite content wire. The grain boundary microstructures were observed by electronback scattering diffraction (EBSD) technique. EBSD was measured with Hitachi S-4300 FE-SEM, TSL solutions camera control system VIT1000, image processing system DSP 2000, and interface controller MSC 2000. The EBSD pattern was analyzed using OIM-Analysis software provided by TSL, Co. Ltd. [24]. Acceleration voltage of SEM beam for the EBSD measurement is 25 kV and the beam current is 15 lA. The surface for EBSD measurements was finally finished by polishing with 0.3 lm deagglomerated alpha alumina powder/ water mixture followed by electropolishing using HClO4 + C2H5OH electrolyte in order to get relatively smooth surface free from surface hardening caused by the mechanical polishing.

Block C

Phase-2 SCC test

15 1.0

Block B for side surface

15 3

3

18

Fig. 2. Schematic of the test and analysis procedures of two HAZ specimens.

numbers (FN) of approximately 13 and 10, respectively. The welds were prepared in accordance with nuclear specification by Gas Tungsten Arc welding. Specimens WOH1 and HTH2 were fabricated by electron-beam welding of the blocks with weld heataffected zones to support blocks to have enough size for preparing crack growth rate specimens. Two contoured double cantilever beam specimens were prepared. Specimen HTH2 was prepared from a weld prepared with a high ferrite content wire, and speci-

2.2. Stress corrosion cracking growth tests The objective of the present work is to quantify SCC growth behavior in 316NG HAZs at different distance from the weld fusion

(a) 300

Vickers hardness, HV0.25

WOH1, P=2.45 N, t=15 s

275 250 225 200 Y= 2.0 mm; Y=6 mm, Y=10mm, Y=12 mm

175

Y=4 Y=8 mm Y=10 mm

150 0

5

10 15 Distance, x (mm)

20

25

(b) Fig. 3. Distribution of Vickers hardness in specimen WOH1. (a) Locations for the HV measurements and (b) HV vs. the distance from the fusion line. Load = 2.45 N, holding time = 15 s.

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(a)

Vickers hardness, HV0.25

300

HTH2, P=2.45 N, t=15 s Y, mm 6 7 8 9 10 11 12 13 14

275 250 225 200 175 150 0

5

10

15 20 Distance, x (mm)

25

30

(b) Fig. 4. Distribution of Vickers hardness in specimen HTH2. (a) Locations for the HV measurements and (b) HV vs. the distance from the fusion line. Load = 2.45 N, holding time = 15 s.

Fig. 5. Vickers hardness distribution in a 316NG HAZ in a weld prepared with a high ferrite content wire [26].

line and to correlate with the microstructural properties and local deformation. SCC tests were performed with two specimens in two phases, i.e., phase-1 SCC and phase-2 SCC, which were designed to

Fig. 6. (a) Location of the crack and (b) SEM morphology of the cracking path on the side surface of specimen WOH1 after phase-1 and phase-2 SCC tests.

measure SCC growth rates in HAZs relatively far from the fusion line or near the fusion line. At first, two 316NG HAZ CDCB specimens (WOH1 and HTH2) of 20 mm thickness were fabricated. The specimen was first air-fatigued under sine wave loading at 15 Hz at Kmax of about 17–18 MPa m0.5 and a load ratio (R = Kmin/Kmax) R = 0.2. After pre-cracking in air, the specimens were side-grooved 5% on each side. Notch tips in both specimens are in the HAZ. The notch tip of specimen WOH1 (after pre-crack in air) is relatively near the weld fusion line. The notch tip of the specimen HTH2 (after pre-cracking in air) is relatively far from the weld fusion line. The specimens were subjected to the phase1 SCC test (including in situ pre-cracking) in high temperature water, then the autoclave was shutdown and two specimens were taken out. One slice was taken from one side of each specimen and the side groove on another side of the specimen was removed to observe the cracking path and crack tip position, Fig. 1. After that, the specimens of 15 mm thickness were pre-cracked (second precracking after the phase-1 SCC test) in air again to make the crack tip near the weld fusion line. After the phase-1 SCC test, the crack tip of specimen WOH1 has already been close to the fusion line. The again-pre-cracked specimens were tested in high temperature water again (phase-2 SCC test). The whole test procedures are summarized in Table 2 and Fig. 2. The in situ pre-cracking in high temperature water under low frequency cyclic loading was designed to facilitate the transition from transgranular cracking

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Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

produced by pre-cracking in air to intergranular stress corrosion cracking in high temperature water. The values of stress intensity factor at different locations of the specimens are calculated with finite element methods (FEM) [25] based on the measurements of the crack length and shape after the SCC tests, which are included in Table 2. FEM is an effective tool for quantifying the crack tip mechanics for cracks of complex geometries. The in situ pre-cracking and stress corrosion crack growth tests were performed in high temperature pure water in an autoclave equipped with a re-circulation system. Pressure was 8.9 MPa, and the flow rate was around 5 L/h. The water chemistry in the re-circulation loop was adjusted to the desired condition then the autoclave temperature was increased to 288 °C. The outlet conductivity was held to be less than 0.16 ls/cm during the test. After reaching the desired water chemistry and then immersing for more than 2 days, the specimens were subjected to in situ fatigue pre-cracking in high temperature water under a series of triangle wave loading. The triangle (Tri.) wave loading mode for in situ pre-cracking is f  0.01 Hz, R = 0.3 (288 cycles, step PMD01), R = 0.5 (576 cycles, step PMD02), and R = 0.7 (864 cycles, step PMD03), respectively. SCC growth tests were performed under constant loading in oxygenated high temperature water at 288 °C with two dissolved oxygen (DO) concentrations: 2 ppm (by weight) and 0.2 ppm. Crack growth was monitored by using an alternating current potential drop (ACPD) machine. The electrode potentials of specimens were measured against a pressure-balanced external Ag–AgCl (0.1 M KCl) reference electrode. After finishing the SCC test, specimens were taken out of the autoclave. Slices were taken at a distance of about 3 mm from the side surfaces of tested CDCB specimens. After slicing, the distribution of hardness was measured using Shimazu Vickers hardness tester with a loading of 2.45 N and a holding time of 15 s. Side surfaces of slices were used for analyzing the cracking paths, microstructure and misorientation in the HAZs and in the crack tip regions. The specimens were then fractured by post-test fatiguing in air. The fracture surfaces were observed with optical microscopy and scanning electron microscope (SEM). The values of the stress corrosion crack length were measured at more than 25 locations along the specimen thickness direction. The average values of SCC crack lengths on each specimen were used to calibrate the crack-growth data obtained from ACPD monitoring. The boundary lines in Fig. 1(a) and (b) are weld fusion lines. Fig. 1 shows that, after the phase-1 SCC test, the crack tip region in specimen WOH1 located at about 1 mm away from the weld fusion line and showed significant branching and intergranular SCC characteristics. The crack tip region in specimen HTH2 is about 3 mm away from the weld fusion line. The side surfaces of two specimens after finishing all test steps are shown in the next section.

metal. More complete Vickers hardness (HV1) values of a HAZ in a weld welded with a high ferrite content wire are shown in Fig. 5 [26]. Generally, hardness is high near the fusion line and tends to decrease with the distance from the fusion line. Optical and SEM photos of the side surfaces of slices cut from specimens after phase-1 and phase-2 SCC tests are shown in Fig. 6 for specimen WOH1 and in Fig. 7 for specimen HTH2. The locations of the cracks in the welds are shown in Figs. 6(a) and 7(a). The shapes of cracks are shown in Figs. 6(b) and 7(b). Extensive crack branching is observed at the crack tip region in specimen WOH1, which shows typical intergranular SCC characteristics, Fig. 6(b). The SCC region in specimen WOH1 is close to the weld fusion line. The starting position for SCC is about 2 mm away from the fusion line and the closest crack tip is about 0.5 mm away from the fusion line. On the side surface, no significant gap could be recognized between SCC in the phase-1 test and that in phase-2 SCC test from the optical observation on the side surface. There is a trend that stress corrosion crack grew toward the fusion line. Two SCC regions are identified on the side surface of specimen HTH2, which correspond to phase-1 SCC and phase-2 SCC, respectively. The phase-1 SCC region is in the HAZ, which is relatively far from the fusion line. The phase-2 SCC tip in specimen HTH2 has penetrated through the weld fusion line and entered the weld metal, Fig. 7(b).

3. Results and discussion 3.1. Hardness distribution and locations of stress corrosion cracks Values of Vickers hardness were measured at different locations on the side surfaces of central blocks of specimens WOH1 and HTH2 after removing the side slices after the phase-1 and phase-2 SCC tests. The locations for the Vickers hardness measurements are shown in Fig. 3(a) and 4(a). The measured values of Vickers hardness at different locations of side surfaces of two tested specimens are shown in Figs. 3(b) and 4(b). Since the cracks are in the regions of hardness measurement, some hardness values near the cracks would be affected by the crack path itself. There is a region of high hardness in the HAZs in both specimens WOH1 and HTH2. The Vickers hardness for the SCC region in specimen WOH1 is about 250 (HV0.25). The Vickers hardness for the phase-1 SCC region in specimen HTH2 is about 235 (HV0.25). Phase-2 SCC tip region is in the weld

Fig. 7. (a) Optical morphology and (b) SEM morphology of the cracking path on the side surface of specimen HTH2 after phase-1 and phase-2 SCC tests.

Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

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Fig. 8. EBSD results (step size = 3 lm) of the crack tip area on the side surface of specimen WOH1. (a) Image, (b) IPF, (c) types of boundaries, and (d) KAM.

3.2. Grain boundary microstructure in HAZs and SCC tip regions Grain boundary microstructural features of the regions near the weld fusion line and SCC regions on the side surfaces of two tested specimens WOH1 and HTH2 were measured by EBSD at locations with different distances from the fusion line. The image for the measured area, measured EBSD patterns in terms of inverse pole

figure, types of grain boundaries such as high angle boundaries and CSLBs, kernel average misorientation (KAM), and summarized data are shown in Figs. 8–12 for specimen WOH1 and in Figs. 13–21 for specimen HTH2. Figs. 6(a) and 8(a) show that there are several intergranular crack branches in the crack tip region of specimen WOH1. The longest branch, Wa, is deviating from the specimen notch direction

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No significant indication of cracking at R3 boundaries can be found. Fig. 8d indicates that KAM in the HAZ near the weld fusion line is higher than that in the region far from the fusion line. Fig. 10 lists the summarized KAM data in Fig. 8, which shows that the fraction of low KAM (0–1) increases and the fraction of high KAM (1–2, and 2–3) decreases with the distance from the weld fusion line, showing high residual strain in the region near the fusion line. Figs. 11 and 12 show the detailed EBSD microstructural characteristics of crack tip region Wa in specimen WOH1. The cracking path along the high angle boundaries is revealed. Relatively high KAM in the

o

0.40 Specimen WOH1

0.35 Fraction of Σ 3 boundaries

Fraction of high angle boundaries (15-180 )

and tends to approach the weld fusion line. The intergranular feature of stress corrosion cracking paths can be seen clearly in Fig. 8(b) and (c). Fig. 8(c) shows a tendency that the ratio of R3 boundaries in the HAZ region adjacent to the fusion line is lower than that in the region relatively far from the fusion line, which is shown more clearly in Fig. 9(a). The fraction of R3 boundaries is lower than 0.1 in the region close to the fusion line, which is more than 0.3 in the regions at relatively long distances from the fusion line, for example, about 20 mm. Fig. 8c shows that stress corrosion cracking proceeds along high angle grain boundaries.

0.30 0.25 0.20 0.15 0.10 0.05 0.00 0

5

10

15

20

25

0.8 WOH1

0.7 0.6 0.5 0.4 0.3 0.2 0

5

10

15

20

25

Distance from the fusion line, DF(mm)

Distance from the fusion line, DF(mm)

(a)

(b)

Fig. 9. (a) Fraction of R3 boundaries vs. distance and (b) fraction of HABs vs. distance from the weld fusion line in specimen WOH1. Measured locations are shown in Fig. 8.

1.0

Specimen WOH1 KAM: 1-2

0.25

Total fraction

Total fraction

0.30

Specimen WOH1 KAM: 0-1

0.8

0.6

0.20 0.15 0.10 0.05 0.00

0.4 0

5

10

15

20

0

25

Distance from the fusion line, DF(mm)

5

10

15

20

25

Distance from the fusion line, DF(mm)

(b)

(a) 0.05 Specimen WOH1 KAM: 2-3

Total fraction

0.04 0.03 0.02 0.01 0.00 0

5 10 15 20 Distance from the fusion line, DF(mm)

25

(c) Fig. 10. KAM vs. distance from the weld fusion line for WOH1 specimen. EBSD step = 2 lm. (a) KAM = 0–1, (b) KAM = 1–2, and (c) KAM = 2–3. Measured locations are shown in Fig. 8.

Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

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Fig. 11. EBSD results (step size = 2 lm) of SCC tip region Wa on the side surface of specimen WOH1. (a) Image, (b) IPF, (c) types of boundaries, and (d) KAM.

Fig. 12. EBSD morphology (step size = 2 lm) of crack tip region Wa on the side surface of specimen WOH1, showing the fusion boundary and the microstructure ahead of the SCC tip. (a) Image, (b) IPF, (c) types of boundaries, and (d) KAM.

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Fig. 13. EBSD results (step size = 3 lm) of the crack tip area including phase-1 SCC region on the side surface of specimen HTH2. (a) Image, (b) IPF, (c) types of boundaries, and (d) KAM.

regions adjacent to the crack tip is shown in Figs. 11 and 12, showing that local strain-hardened area acts as preferential stress corrosion cracking path.

Fig. 13 shows the EBSD results of regions on the side surface of HTH2 specimen, which includes the IGSCC tip occurred during the phase-1 SCC period. Fractions of different types of boundaries are

1925

o

Ratio of high angle grain boundaries (15-180 )

Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

Σ 3 grain boundaries

Ratio of Σ 3 grain boundaries

0.40

Specimen HTH2 (PI-SCC region)

0.35 0.30 0.25 0.20 0.15 0.10 0.05 0

5

10

15

20

25

Distance from the fusion line, DF(mm)

1.0 Specimen HTH2 (PI-SCC region)

0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0

5 10 15 20 Distance from the fusion line, DF(mm)

(a)

25

(b)

Fig. 14. (a) Fraction of R3 boundaries and (b) ratio of high angle grain boundaries (15–180°) distribution as functions of the distance from the weld fusion line for specimen HTH2 in phase-1 SCC region. Measured locations are shown in Fig. 13.

0.25

Specimen HTH2 (P1-SCC) KAM: 0-1

1.00

Specimen HTH2 (P1-SCC) KAM: 1-2

0.20 Total fraction

Total fraction

0.95 0.90 0.85 0.80 0.75

0.15

0.10

0.05

0.70 0

5

10

15

20

0

25

5

10

15

20

25

Distance from the fusion line, DF(mm)

Distance from the fusion line, DF(mm)

(b)

(a) 0.04

KAM2To3

Specimen HTH2 (PI-SCC) KAM: 2-3

Total fraction

0.03

0.02

0.01

0.00 0

5

10

15

20

25

Distance from the fusion line, DF(mm)

(c) Fig. 15. KAM distribution as function of the distance from the weld fusion line for specimen HTH2 in phase-1 SCC region. (a) KAM = 0–1, (b) KAM = 1–2, and (c) KAM = 2–3. Measured locations are shown in Fig. 13.

shown in Fig. 14. KAM data are summarized in Fig. 15. Figs. 14 and 15 show that changes in the grain boundary characteristics, such as increasing fraction of R3 boundaries, increasing fraction of high angle boundaries and decreasing misorientation as indicated by KAM with increasing distance from the weld fusion line. The results are consistent with the results obtained on specimen WOH1. Figs. 16 and 17 show the EBSD results of typical IGSCC

regions occurred during the phase-1 SCC test, where typical intergranular cracking along HAB is identified. Fig. 17 also shows one microstructure configuration as preferential stress corrosion cracking path: the random boundary adjacent to high density R3 boundaries is likely to be sensitive to stress corrosion cracking, which has been observed in previous results for SCC of coldworked 316NG stainless steel in simulated BWR environments

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Fig. 16. EBSD (step size = 3 lm) results of the side surface of specimen HTH2. (a) Image, (b) IPF, (c) types of boundaries, and (d) KAM.

Fig. 17. EBSD results (step size = 1 lm) of one part of the phase-1 SCC region on the side surface of specimen HTH2. (a) Image, (b) IPF, and (c) types of boundaries.

[21]. Local deformation as the result of the effect of prior deformation on alloys with microstructural heterogeneity would play an important role in the active SCC path. Fig. 18 shows the EBSD results of phase-2 SCC area and nearby regions on the side surface of HTH2 specimen. Fractions of different types of boundaries are shown in Fig. 19 and KAM data are

summarized in Fig. 20. The EBSD results are consistent with the results in Figs. 14 and 15 on another side surface area of specimen HTH2 and the results in Figs. 9 and 10. Fraction of R3 boundaries increases, fraction of high angle boundaries increases, and misorientation represented by KAM decreases with increasing the distance from the weld fusion line. EBSD results measured on total

Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

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Fig. 18. EBSD (step size = 3 lm) results of the side surface of specimen HTH2 of phase-2 SCC tip and nearby area. (a) Image, (b) IPF, (c) types of boundaries, and (d) KAM.

three locations on two specimens show consistent results on the grain boundary properties and misorientation distribution.

Fig. 21 shows the detailed EBSD analysis of phase-2 SCC tip region. Phase-2 stress corrosion crack has penetrated into the weld metal.

Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

Region H4

0.85

Total fraction

Region H3 Region H1

0.90

0.15

Region H2

0.95 Total fraction

0.20

Specimen HTH2 KAM: 0-1

Region H1

1.00

0.10

0.05

0.80

Specimen HTH2 KAM: 1-2 Region H2 Region H3 Region H4

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0.00 0

2

4

6

8

10

12

14

16

18

20

0

Distance from the fusion line, DF(mm)

2

4

6

8

10

12

14

16

18

20

Distance from the fusion line, DF(mm)

(a)

(b)

o

0.35 Specimen HTH2

Ratio of Σ 3 boundaries

Ratio of high angle boundaries (15-180 )

Fig. 19. KAM vs. distance from the weld fusion line for HTH2 specimen. EBSD step = 3 lm. (a) KAM = 0–1, (b) KAM = 1–2. The measured regions are shown in Fig. 18.

0.30 0.25 0.20 0.15 0.10 0

2

4

6

8

10

12

14

16

18

20

Distance from the fusion line, DF(mm)

(a)

0.8 Specimen HTH2

0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0

2

4

6

8

10

12

14

16

18

20

Distance from the fusion line, DF(mm)

(b)

Fig. 20. (a) Fractions of R3 boundaries and (b) fraction of high angle (15–180°) boundaries in the regions with phase-2 SCC area on the side surface of specimen HTH2 as function of the distance from the weld fusion line, based on the EBSD (step size = 3 lm) results. The measured regions are shown in Fig. 18.

The cracking path inside the weld metal is close to the dendrite direction of the weld metal. 3.3. Fracture surfaces and stress corrosion cracking growth kinetics The optical photos of the fracture surfaces of two specimens WOH1 and HTH2 are shown in Fig. 22. Typical SEM morphologies of the fracture surfaces are shown in Figs. 23 and 24. Typical IGSCC on the fracture surface of specimen WOH1 is identified in Fig. 22(a). The boundary between phase-1 SCC and phase-2 SCC regions can not be easily identified in the optical photo, which can be identified based on the SEM photo in Fig. 23. Two SCC regions on the fracture surface of specimen HTH2 can be identified in Fig. 22(b). Typical IGSCC feature is observed in the phase-1 SCC region, which is shown more clearly in Fig. 24(a). Phase-2 SCC region is deviating from the normal plane, showing kinking features with partly intergranular cracking facets, Fig. 24(b). The values of stress corrosion crack length were measured at more than 25 locations along the specimen thickness direction. The average values of crack lengths on each specimen were used to calibrate the crack-growth data obtained from ACPD monitoring. The crack lengths in phase-1 and phase-2 SCC regions for specimen WOH1 in the total crack length are divided via selecting typical SEM photos with distin-

guishable boundary to calculate the ratio between crack advance in phase-1 SCC and that in phase-2 SCC. The calibrated cracklengths vs. test time from ACPD monitoring are shown in Figs. 25–28. In the SCC step C11 after in situ pre-cracking for specimen WOH1, the crack grew slowly during the first 250 h. As deduced from Fig. 25, the crack growth rate in 2 ppm DO water increased with testing time until it reached a quasi-steady growth state of 4.55  1010 m/s in the test time period of 600–900 h. Crack growth rate in 2 ppm DO water was about 3.90  1010 m/s in the test time period of 900–1157 h, which was somehow lower than that in the period of 650–900 h. Crack growth rate was about 2.51  1010 m/s in the step C12 in phase-1 SCC test for specimen WOH1 tested in 0.2 ppm DO water. The incubation period in the SCC step C21 was about 100 h, which was significantly shorter than that in the step C11. Short SCC incubation period in step C21 was probably caused by the fact that the crack front at the starting of step C21 was more close to intergranular type thus the transition to intergranular SCC was easier and took less time than that in the step C11. The quasi-steady state CGR in step C21 for WOH1 is 1.01  1010 m/s in the period of 300–1434 h. In the SCC step C11 after in situ pre-cracking for specimen HTH2, the crack grew slowly during the first 700 h. As deduced

Z. Lu et al. / Corrosion Science 53 (2011) 1916–1932

1929

Fig. 21. EBSD (step = 3 lm) results in terms of (a) image, (b) IPF, (c) types of boundaries, and (d) KAM for phase-2 SCC tip area on the side surface of specimen HTH2.

from Fig. 27, the crack growth rate increased with testing time until it reached a quasi-steady growth state with a CGR of 2.15  1010 m/s in the test time period of 900–1157 h. After changing DO from 2 to 0.2 ppm, CGR remained almost unchanged in the initial 40 h period (1157–1200 h), then changed to a new quasi-steady state with a CGR of 6.95  1011 m/s in the following period of about 100 h. In the step C21 in phase-2 SCC test for specimen HTH2, the crack grew slowly during the first 320 h. As deduced from Fig. 28, the crack growth rate increased with testing time until it reached a quasi-steady growth state of 1.23  1010 m/s in the test period of 450–1434 h. Stress corrosion cracking growth incubation periods in the SCC test steps after in situ pre-cracking are shown in Figs. 25, 27 and 28, which has been observed in previous experiments for coldrolled or warm-rolled low-carbon stainless steels and 316NG weld HAZs and weld metals [10–13,20]. Among several possible causes for such a SCC incubation period, the transition from transgranular (TG) cracking during in situ pre-cracking to intergranular (IG) cracking during SCC would play a dominant role in the incubation period before the onset of steady crack growth. The incubation period in the step C21 for specimen WOH1 is less significant than that in step C12 and those in test steps C11 and C21 for specimen HTH2. Previous results on SCC growth of warm-rolled 304L SS in the longitudinal–transverse (L–T) orientation in 288 °C pure water showed that SCC incubation period was almost not observed if the TG–IG transition was not significant even after long time SCC testing. These experimental observations are consistent with the proposed IG–TG transition theory for the SCC incubation periods after in situ pre-cracking [10–13,20,21]. Crack growth rates in the test step C11 in specimen WOH1 are about two times higher than that in specimen HTH2 under the same testing condition. The difference is the location of the cracking area in the HAZ and the resultant differences in microstructure and degree of strain-hardening represented by hardness and mis-

orientation. The distance between phase-1 SCC region and the weld fusion line in specimen WOH1 is shorter than that in specimen HTH2, Figs. 6 and 7. Figs. 3–5 show that Vickers hardness in the weld HAZ generally decreases with increasing the distance from the weld fusion line. Similar results for hardness distribution in 316NG HAZ from PLR piping mockup have been reported [27]. Experimental data show that crack growth rate of 316NG HAZ in simulated BWR pure water increases with increasing hardness [27]. Detailed microstructural analyses show that, in the weld HAZ, kernel average misorientation decreases and ratio of R3 boundaries increases with increasing the distance from the weld fusion line, as shown in Figs. 8(d), 13(d), 18(d), and 10, 15, and 19. All the results show that a high degree of strain-hardening leads to a high stress corrosion cracking growth rate. A summary of crack growth rates in oxygenated 288 °C pure water for non-sensitized stainless steels at different yield strengths is shown in Fig. 29, based on the published data [9–13,19,28–31], which clearly shows that crack growth rate increases with increasing yield strength as the results of strain-hardening. The effect of strainhardening in terms of yield strength or hardness on SCC growth rate in high temperature water have been reported and analyzed [1,6,28,30,32–39]. SCC growth rate in high temperature water is formulated based on interaction between crack tip mechanics and crack tip oxidation kinetics [18,34,35,40]. Two necessary conditions for the progressing of SCC, localization of oxidation and acceleration of oxidation, are proposed by Shoji [18]. Yield strength would have potential effects on oxidation rate constants and the distance at which the crack tip oxidation and mechanics interacts [34]. High yield strength would enhance the crack tip oxidation within localized and constrained crack tip area due to highly localized stress and strain/strain rate, and possibly due to the high dislocation density in the strain-hardened alloys. The role of R3 boundaries needs more systematic investigation since observed cracking paths generally follow random grain boundaries.

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Steady state crack growth rate for specimen WOH1 in the phase-1 SCC test in 288 °C pure water with 2 ppm DO is about 70% higher than that in 288 °C pure water with 0.2 ppm DO. Steady state crack growth rate for specimen HTH2 in the phase-1 SCC test in 288 °C pure water with 2 ppm DO is about three times higher than that in 288 °C pure water with 0.2 ppm DO. These results imply that the effect of dissolved oxygen or electrochemical potential is more significant for non-hardened or less-hardened alloys than that for heavily hardened alloys [6,21,41]. It has been reported that the SCC mitigation factor (SCCMF) by decreasing electrode potential is less significant for cold-worked stainless steels and cold worked than that for non-hardened stainless steels, Eq. (1) and Table 3 [6,29,40–42]. SCCMF is also found to be strongly dependent on material chemistry. For example, Andresen and Morra reported that the effect of electrode potential is minor for rolled alloys with high content of Si [42].

SCCMF ¼ CGRðECP1Þ=CGRðECP2Þ

ð1Þ

SCCMF is the SCC mitigation factor. CGR(ECP1) and CGR(ECP2) are crack growth rates at electrode potentials ECP1 and ECP2, respectively. The details of the transition growth periods after changing DO from 2 ppm DO to 0.2 ppm DO are different in specimens WOH1 and HTH2, which are possibly caused by the crack tip geometry and related change of local water chemistry and electrochemistry. Fig. 23. A typical SCC region on the fracture surface of specimen WOH1.

Fig. 22. The fracture surfaces of specimens: (a) specimen WOH1 and (b) specimen HTH2.

Fig. 24. Typical fracture surfaces of HTH2 specimen. (a) Phase-1 SCC typical and (b) phase-2 SCC typical.

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0.5 Step C11 2 ppm DO

1.2 1.0

Step C12 0.2 ppm DO

0.8 0.6 316L HAZ, WOH1 phase-1 SCC 0.4 CDCB specimen 0 Pure water, 288 C 0.2 Constant loading

2 ppm DO 0.2 ppm DO

Average crack advance, Δ a(mm)

Average crack advance, Δ a(mm)

1.4

200

400

0.2 0.1

200

1E-8

0.4

Crack growth rate, CGR(m/s)

316L HAZ, WOH1 phase-2 SCC CDCB specimen 0 Pure water, 288 C, 2 ppm DO Constant loading

0.3

0.2

0.1

400

600 800 1000 1200 1400 1600 Test time, t(h)

Fig. 28. Crack advance vs. test time for the 316NG HAZ specimen HTH2 during the phase-2 SCC test in 288 °C pure water with 2 ppm DO.

0.5 Average crack advance, Δ a(mm)

0.3

0

600 800 1000 1200 1400 1600 Test time, t(h)

Fig. 25. Crack advance vs. test time for the 316NG HAZ specimen WOH1 during phase-1 SCC test in pure water at 288 °C.

0.4

0.0

0.0 0

316L HAZ, HTH2 phase-2 SCC, CDCB specimen 0 Pure water, 288 C, 2 ppm DO Constant loading

Andresen et al, SA & CW 304L, 2 ppm DO Andresen et al, CW 316L, 2 ppm DO FRI, CW 304L, 2 ppm DO FRI, CW 316L, 2 & 7.5 ppm DO Tsubota et al., SA & CW 316L, 8 ppm DO (YS converted) Itow et al., 304L, 12-20 ppm DO Itow et al., 316L, 12-22 ppm DO Itow et al., 316NG, 12-20 ppm DO

1E-9

1E-10

1E-11 o

Low-C SS in oxygenated 288 C pure water 0.5 Fracture mechanics specimens, K:25-35MPa.m

o

YS at ~288 C

0.0

0

200

400

600 800 1000 1200 1400 1600 Test time, t(h)

Fig. 26. Crack advance vs. test time for the 316NG HAZ specimen WOH1 during phase-2 SCC test in 288 °C pure water with 2 ppm DO.

200

300

400

500

600

700

800

Yield strength, YS(MPa) Fig. 29. The effect of yield strength on the SCC growth rate of non-sensitized stainless steels in oxygenated 288 °C pure water [9–13,28–31].

Table 3 Calculated values of SCCMF based on the experimental data in Refs. [6,29,40–42] for various alloys in simulated BWR environments.

0.5 Average crack advance, Δ a(mm)

1E-12 100

Step C11 2 ppm DO

0.4

Step C12 0.2 ppm DO

0.3 0.2

316L HAZ, HTH2 phase-1 SCC CDCB specimen 0 Pure water, 288 C Constant loading

0.1

step C11, 2 ppm DO step C12, 0.2 ppm DO

0.0 0

200

400

600 800 1000 1200 1400 1600 Test time,t(h)

Fig. 27. Crack advance vs. test time for the 316NG HAZ specimen HTH2 during the phase-1 SCC test in 288 °C pure water.

Types of alloys

Sensitized 304 SS

Solutionannealed low-C SS

CW low-C SS and alloy 600

Warm-rolled 5Si15Ni12Cr15

ECP1, mV(SHE) ECP2, mV(SHE) SCCMF = CGR(ECP1)/ CGR(ECP2)

150 500 >20

160 230 10

550 150 3

550 150 Close to 1

incubation period, exhibiting a similar steady state CGR to that in specimen WOH1 tested at the same time, even though the final crack front of specimen HTH2 has been inside the weld metal. The crack growth behavior of 316NG weld metal in simulated BWR environments has been reported [7,12,23,43,44].

4. Conclusions Steady state crack growth rate in the step C21 in phase-2 SCC test for specimen WOH1 is about one-fourth of that in the step C11 in phase-1 SCC test, which could be caused by the de-activation of some branching cracks after the second pre-cracking in air, and the changing of the local microstructure and possibly residual stress distribution. In the step C21 in the phase-2 SCC test for specimen HTH2, steady state crack growth is observed after the

Stress corrosion crack growth in 316NG weld in three regions: in the HAZ relatively far from the fusion line, in the HAZ relatively close to the fusion line and in the weld metal near the fusion line in oxygenated pure water at 288 °C was investigated. EBSD results show that, with increasing the distance from the fusion line, kernel average misorientation decreases and fraction of R3 boundaries in-

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creases. SCC proceeds via intergranular mode and cracking path has been found to follow random grain boundaries rather than R3 boundaries. Decreasing dissolved oxygen concentration reduces crack growth rate, which is more effective in the HAZ with a lower hardness than that in the HAZ with a higher hardness. EBSD results reveal that local deformation as the result of submicrostructure produced by cold work or welding process would play an important role in determining the stress corrosion cracking path and kinetics. Higher SCC susceptibility was revealed in the HAZ than in the base metal related to grain boundary properties and local strain-hardening. Acknowledgments

[19]

[20] [21] [22]

[23] [24] [25]

This work has been performed as a part of the PEACE-E program jointly supported by EDF, EPRI, SSM, TEPCO, KEPCO, TohokuEPCO, ChubuEPCO, JAPCO, HITACHI Ltd., MHI, TOSHIBA Co., and IHI. This work has been also performed under the support of Grant-in-Aid for Scientific Research (S) 17106002 and (C) 20560063, Japan Society for the Promotion of Science, and the support of the International Cooperative Program for Education and Research, the Japanese Ministry of Education, Culture, Sports, Science and Technology. Part of this work was performed as a part of the contracted work on Enhancement of Ageing Management and Maintenance of Nuclear Power Plants by Nuclear and Industrial Safety Agency (NISA) in the Ministry of Economy, Trade and Industry (METI) of Japan. F.J. Meng thanks the support of The Special Funds for the Major State Basic Research Projects G2006CB605000 (973 Program, China). References [1] T. Shoji, Progress in the mechanistic understanding of BWR SCC and its implication to the prediction of SCC growth behavior in plants, in: Proceedings of the 11th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, Skamania, Stevenson, Washington, ANS, 2003, pp. 588–598. [2] S. Suzuki, K. Kumagayi, C. Shitara, J. Mizutani, A. Sakashita, H. Tohkuma, H. Yamashita, Maintenology 3 (2004) 65–70. [3] S. Suzuki, K. Takamori, K. Kumagayi, A. Sakashita, N. Yamashita, C. Shitara, Y. Okamura, E-J. Adv. Maint. 1 (2009) 1–29. [4] NISA, Documents Presented at Fifth Structural Integrity Evaluation Committee, February 18, 2003. [5] T.M. Angeliu, P.L. Andresen, E. Hall, J.A. Sutliff, S. Sitzman, R.M. Horn, Intergranular stress corrosion cracking of unsensitized stainless steels in BWR environments, in: Proceedings of the Ninth International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, TMS, 1999, pp. 311–317. [6] P.L. Andresen, M.M. Morra, J. Nucl. Mater. 383 (2008) 97–111. [7] M. Itow, S. Suzuki, SCC growth and retardation behavior of low carbon stainless steel piping in the region near the weld fusion boundary, in: 55th Material and Corrosion Symposium, Paper F08104, JSCE, 2008, pp. 31–34. [8] T. Arai, K. Watanabe, K. Kako, Y. Miyahara, SCC growth behavior of stainless steel in the region near the weld fusion boundary, in: 55th Material and Corrosion Symposium, Paper F08107, JSCE, 2008, pp. 35–38. [9] P.L. Andresen, T.M. Angeliu, L.M. Young, Effect of Martensite and hydrogen on SCC of stainless steels and alloy 600, in: Proceeding of the Corrosion/2001, NACE, Paper No. 228, 2001. [10] Z.P. Lu, T. Shoji, Y. Takeda, A. Kai, Y. Ito, Corrosion 63 (2007) 1021–1032. [11] Z.P. Lu, T. Shoji, Y. Takeda, Y. Ito, S. Yamazaki, Corros. Sci. 50 (2008) 698–712. [12] Z.P. Lu, T. Shoji, Y. Takeda, Y. Ito, A. Kai, N. Tsuchiya, Corros. Sci. 50 (2008) 625– 638. [13] Z.P. Lu, T. Shoji, Y. Takeda, Y. Ito, A. Kai, S. Yamazaki, Corros. Sci. 50 (2008) 561– 575. [14] Z. Jiao, G.S. Was, J. Nucl. Mater. 382 (2008) 203–209. [15] B. Alexandreanu, B.H. Sencer, V. Thaveeprungsriporn, G.S. Was, Acta Mater. 51 (2003) 3831–3848. [16] B. Alexandreanu, G.S. Was, Corrosion 59 (2003) 705–720. [17] S. Lozano-Perez, T. Yamada, T. Terachi, M. Schroder, C.A. English, G.D.W. Smith, C.R.M. Grovenor, B.L. Eyre, Acta Mater. 57 (2009) 5361–5381. [18] T. Shoji, Localized and accelerated oxidation and stress corrosion cracking-role of stress, strain, hydrogen and microstructures, in: Proceedings of the

[26]

[27]

[28]

[29]

[30] [31]

[32] [33]

[34] [35]

[36]

[37] [38] [39] [40] [41]

[42]

[43]

[44]

Quantitative Micro–Nano (QMN) Approach to Predicting Stress Corrosion Cracking in Water-Cooled Nuclear Plants, 2010, CDROM. P.L. Andresen, L.M. Young, W.R. Catlin, R.M. Horn, Stress corrosion crack growth rate behavior of various grades of cold worked stainless steel in high temperature water, in: Proceedings of the Corrosion/2002, NACE, Paper No. 511, 2002, CDROM. Z.P. Lu, T. Shoji, T.C. Dan, Y.B. Qiu, T. Yonezawa, Corros. Sci. 52 (2010) 2547– 2555. Z.P. Lu, T. Shoji, F.J. Meng, Y.B. Qiu, T.C. Dan, H. Xue, Corros. Sci. 53 (2011) 247– 262. Z.P. Lu, K. Sakaguchi, K. Negishi, Y. Takeda, Y. Ito, T. Shoji, Quantifying the effects of strain-hardening and water chemistry on crack growth rates of 316L SS welds in high temperature water, in: Proceedings of the 14th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, ANS, 2009, pp. 636–645, CDROM. J.H. Kim, R.G. Ballinger, Corrosion 64 (2008) 645–656. OIM Analysis for Windows 5.0, User Manual, TexSEM Laboratories, Inc., Utah, USA, 2005. Z.P. Lu, H. Xue, T. Shoji, Crack branching and its effect on environmentally assisted cracking in high temperature water environments, in: Proceedings of the 2010 ASME Pressure Vessels and Piping Division Conference, ASME, 2010, PVP2010-25818. K. Negishi, Evaluation of stress corrosion cracking characteristic of weldment of 316L stainless steels piping mockups influenced by thermal aging, Master Thesis, Department of Mechanical and Aerospace Engineering, Tohoku University, March 2009. M. Ando, K. Nakata, M. Itow, N. Tanaka, M. Koshiishi, R. Obata, Y. Miwa, Y. Kaji, M. Hayakawa, CGR behavior of low carbon stainless steel of hardened heat affected zone in plr piping weld joints, in: Proceedings of the 13th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, 2007, CNS, CDROM. T. Shoji, G.F. Li, J.H. Kwon, S. Matsushima, Z.P. Lu, Quantification of yield strength effects on IGSCC of austenitic stainless steels in high temperature waters, in: Proceedings of the 11th International Conference on Environmental Degradation of Materials Nuclear Power Systems – Water Reactors, ANS, 2003, pp. 834–843. M. Itow, M. Kikuchi, N. Tanaka, J. Kuniya, M. Yamamoto, S. Suzuki, S. Namatame, T. Futami, SCC growth rates and reference curves for low-carbon stainless steels in BWR environments, in: Proceedings of the 2004 ASME Pressure Vessels and Piping Division Conference, ASME, 2004, pp. 167–173. P.L. Andresen, M.M. Morra, J. Pressure Vessel Technol. 129 (2007) 488–506. M. Tsubota, Y. Katayama, Y. Saito, T. Kaneko, Stress corrosion cracking growth behavior of cold worked austenitic stainless steel in high temperature water, in: Proceedings of the 12th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, TMS, 2005, pp. 167–181. S. Yamazaki, Z.P. Lu, Y. Ito, Y. Takeda, T. Shoji, Corros. Sci. 50 (2008) 835–846. P.M. Scott, P. Combrade, On the mechanism of stress corrosion crack initiation and growth in alloy 600 exposed to PWR primary water, in: Proceedings of the 11th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, ANS, 2003, pp. 29–38. T. Shoji, Z.P. Lu, H. Murakami, Corros. Sci. 52 (2010) 769–779. T. Shoji, S. Suzuki, R.G. Ballinger, Theoretical prediction of SCC growth behavior-threshold and plateau growth rate, in: Proceedings of the Seventh International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, NACE, 1995, pp. 881–891. T. Shoji, Z.P. Lu, S. Yamazaki, The effect of strain-hardening on PWSCC of nickel-base Alloys 600 and 690, in: Proceedings of the 14th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, ANS, 2009, pp. 220–238. K. Arioka, T. Yamada, T. Terachi, Corrosion 62 (2006) 568–575. W.C. Moshier, C.M. Brown, Corrosion 56 (2000) 307–320. Z.P. Lu, Y. Takeda, T. Shoji, J. Nucl. Mater. 383 (2008) 92–96. P.L. Andresen, F.P. Ford, Mater. Sci. Eng. A 103 (1988) 167–184. M. Toloczko, S.M. Bruemmer, P.L. Andresen, SCC crack growth of cold-worked 316L SS in BWR oxidizing and hydrogen water chemistry conditions, in: Proceedings of the 13th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, CNS, 2007, CDROM. P.L. Andresen, M.M. Morra, Effects of Si on SCC of irradiated and unirradiated stainless steels and nickel alloys, in: Proceedings of the 12th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, TMS, 2005, pp. 87–106. K. Nakade, T. Hirasaki, S. Suzuki, K. Takamori, K. Kumagai, Y. Tanaka, K. Umeoka, SCC crack propagation behavior in 316L weld metal under high temperature water, in: JSCE Conference on Materials and Corrosion 2008, Paper No. A-204. Z.P. Lu, T. Shoji, Y. Takeda, Y. Ito, A. Kai, N. Tsuchiya, Effects of electrochemical potential on crack growth rates of 316L weld HAZ and weld metal materials in 288 °C pure water, in: Proceedings of the 13th International Conference on Environmental Degradation Materials Nuclear Power Systems – Water Reactors, CNS, 2007, CDROM.