Stress corrosion cracking in the heat affected zone of a stainless steel 308L-316L weld joint in primary water

Accepted Manuscript Title: Stress Corrosion Cracking in the Heat Affected Zone of a Stainless Steel 308L-316L Weld Joint in Primary Water Author: Lijin Dong Qunjia Peng En-Hou Han Wei Ke Lei Wang PII: DOI: Reference:

S0010-938X(16)30073-7 http://dx.doi.org/doi:10.1016/j.corsci.2016.02.030 CS 6667

To appear in: Received date: Revised date: Accepted date:

4-9-2015 19-12-2015 15-2-2016

Please cite this article as: Lijin Dong, Qunjia Peng, En-Hou Han, Wei Ke, Lei Wang, Stress Corrosion Cracking in the Heat Affected Zone of a Stainless Steel 308L-316L Weld Joint in Primary Water, Corrosion Science http://dx.doi.org/10.1016/j.corsci.2016.02.030 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Stress Corrosion Cracking in the Heat Affected Zone of a Stainless Steel 308L-316L Weld Joint in Primary Water Lijin Donga,b, Qunjia Penga,*, En-Hou Hana, Wei Kea, Lei Wangb a

Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang City 110016, China b

Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education), Northeastern University, Shenyang City 110819, China

* Corresponding author. Tel.: +86 24 2384 1676; Fax: +86 24 2389 4149. E-mail address: [email protected] (Q.J. Peng)

Highlights 

Studied stress corrosion cracking (SCC) in a stainless steel 308L-316L weld joint.

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SCC growth in the heat affected zone (HAZ) affected by chemistry of the primary water.

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SCC grew in the HAZ in off-normal water chemistry with dissolved oxygen (DO).

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No SCC growth in the HAZ in normal water chemistry with dissolved hydrogen.

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Low SCC growth rate in the HAZ likely relates to the low residual strain level.



SCC in the HAZ in primary water with DO follows the slip-oxidation mechanism.

ABSTRACT Stress corrosion cracking (SCC) in the heat affected zone (HAZ) of a stainless steel 308L-316L weld joint in primary water of pressurized water reactor was investigated. Stress corrosion crack growth in the HAZ was observed in off-normal primary water chemistry with dissolved oxygen, but not in normal primary water chemistry with dissolved hydrogen. This suggests that it is unlikely a stress corrosion crack propagating in the HAZ could reach the fusion 1   

boundary and penetrate into the weld metal under normal primary water chemistry conditions. Microstructure analysis of the crack tip suggests that the SCC follows the slip-oxidation mechanism. Keywords: A. Stainless steel; B. SEM; B. TEM; C. Stress corrosion; C. Welding

1. Introduction Austenitic stainless steel has been widely used as a key structural material in the primary circuit of pressurized water reactors (PWRs) due to its high resistance to corrosion and stress corrosion cracking (SCC) in primary water. The steel, however, was found to be susceptible to SCC by both laboratory experiments and field experiences [1-8]. The SCC is typically associated with off-normal PWR primary water chemistry and cold working (CW) of the steel. Dissolved oxygen (DO) and harmful anions such as Cl- and SO42- promote the occurrence of SCC [9-15]. The concentration ratio of B/Li of the primary water also promotes SCC if bubbling oxygen into the water, while it has little effect on SCC in deaerated water [16-17]. In normal PWR primary water chemistry with dissolved hydrogen (DH) at low electrochemical potential conditions, CW is a key factor leading to SCC of the steel [6-7, 18-19]. The promotion of SCC by CW is attributed to the damage at the grain boundaries and the increase in hardness and yield strength [6, 19-22]. Studies have revealed that in the heat affected zone (HAZ) of a weld joint, residual strain as a result of weld shrinkage increases the yield strength, leads to heterogeneous distribution of strain, and therefore, promotes SCC as in cold worked alloys [23-28]. The SCC behavior in the fusion boundary (FB) region of a weld joint has also been investigated. In a weld joint of stainless steels, an investigation revealed that a stress corrosion crack propagating in the HAZ toward the FB in high temperature water with dissolved oxygen tended to avoid the weld metal and showed slower growth after it reached the FB, primarily due to a change in the direction of the intergranular propagation [29]. This suggested that the FB could be a barrier for SCC propagation. A couple of studies on SCC behavior of the FB region of an Alloy 182-low alloy steel (LAS) weld joint were also conducted [12-15]. The studies showed that a stress corrosion crack propagating perpendicular to the FB in the weld metal in high temperature pure water tended to cease propagation after it reached the FB [12-13].

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However, increasing the stress intensity factor, K to be more than 60 MPa√m or adding chloride anion into the water could cause further propagation of SCC from the FB into the LAS [14-15]. A stainless steel 308L-316L weld joint is commonly used to attach the safe end of primary circuit piping to reactor pressure vessel nozzles. To date, few SCC incidences of the weld joint have been reported, primarily due to a high SCC resistance of the 308L weld metal in primary water. Nevertheless, the HAZ in 316L stainless steel in a safe-end weld joint has exhibited SCC susceptibility in primary water [30-31]. This raises a concern for the SCC growth behavior in the HAZ as well as in the FB region. To the author's knowledge, limited experimental data on SCC in the HAZ of 316L stainless steel in primary water has been reported. In the current study, SCC in a 308L-316L stainless steel weld joint in simulated PWR primary water was studied by crack growth rate (CGR) testing and was followed by characterization of the microstructure of the weld joint. The objective is to clarify the growth behavior of a stress corrosion crack propagating in the HAZ to the FB, with focuses on the effect of PWR primary water chemistry and stress intensity factor. 2. Experimental method 2.1 Materials and specimen The weld joint used for the study was cut from a mockup of the safe-end weld joint. The composition of the weld and base metals of the weld joint is listed in Table 1. The weld joint was prepared by multi-pass shielded metal arc welding with no post-weld heat treatment. A 0.5-inch thick, compact tension (0.5T-CT) specimen was used for the CGR test. The specimen was extracted from the weld joint with the orientation shown in Fig. 1. The orientation is employed to simulate the propagation of SCC to the FB in the HAZ, since it was reported that SCC in the HAZ of 316 stainless steel had propagated to the FB at an angle of around 45 in a safe-end weld joint of PWR primary circuit piping [31]. The notch-tip in the specimen is approximately 3.5 mm from the FB in the HAZ. Before the CGR test, precracking in air by fatigue to a crack length increment of approximately 2.1 mm at a maximum stress intensity factor, Kmax of less than 15 MPa√m was conducted. The precrack length was controlled so that the crack did not penetrate the FB. As shown in Fig. 2, the crack tip on the side faces of the specimen was 1.12 mm and 1.22 mm away from the FB, respectively. A 10 mm × 10 mm × 2 mm plate containing the FB was also extracted from the weld joint for microstructure analysis, as shown in Fig. 1. 3   

2.2 Microstructure analysis An optical microscope (OM) was used for the observation of the metallographic microstructure of the weld. An FEI-XL30 field emission scanning electronic microscope (SEM) equipped with a camera that is used for analyzing the misorientation of grain boundaries by electron backscatter diffraction (EBSD) in connection with the TSL software was used for observation and microstructure analysis of the grain boundary and crack propagation path in the HAZ. The specimens for EBSD were mechanically ground by silicon carbide paper up to 2000 grit, then polished using 1 μm diamond paste, and finally polished with a 0.04 μm colloidal silica polishing slurry. The EBSD analysis was performed with a step size of 3 μm at a voltage of 20 keV. Qualitative strains were estimated by measuring the kernel average misorientation (KAM), which was calculated by indexing and averaging the misorientations of neighbor-to-neighbor points within the grains that were obtained by EBSD analysis. For a given point, the average misorientation of that point with all of its neighbors was calculated with a criterion that misorientations exceeding a tolerance value (5º here) were excluded from the calculation. A hardness map of the weld joint was obtained using a HV-1000 Vickers microhardness tester with a load of 100 g and a holding time of 10 s. 2.3 CGR test The CGR test was performed in simulated PWR primary water at 320 °C in a 3-L autoclave with a flow rate of 5 L/h. The primary water was prepared by high-purity water with 1200 mg/L of B as H3BO3 and 2.3 mg/L of Li as LiOH. DO and DH in the influent water was controlled by bubbling either gas mixtures of N2+O2 or H2 into the water tank. Conductivity and DO of the influent water were continuously monitored during the test by InPro 6050 and InPro 7000-VP sensors from Mettler Toledo, while DH was monitored by a DH-35A sensor from DKK-TOA. Crack extension was continuously monitored during the test using a reversing direct current potential drop (DCPD) technique, which has a noise resolution of 1~3 μm. Details of the setup and data collection process of the DCPD system were reported previously [32]. The CT specimen was loaded by clevis pins and electrically isolated from the pins and clevises using zirconia sleeves inserted into the specimen, and zirconia washers inserted into the space between the side faces of the specimen and the clevises. Prior to the SCC test, three steps of triangular waveform loadings for in-situ precracking, and two steps of trapezoidal waveform loadings for 4   

transformation of the crack growth from transgranular mode to intergranular mode were employed in the environment. Details of the loading conditions are listed in Table 2. The stress corrosion crack growth behavior was tested under constant K conditions. To investigate the effects of DO, DH and K on CGR, in total three DO levels at <5, 200, and 2000 μg/L, two DH levels at 870 and 2600 μg/L and four K levels at 25, 30, 35 and 40 MPa√m were employed for the test. During the test, trapezoidal waveform loading was occasionally applied to activate the crack growth. The loading and environment conditions for the SCC test are summarized in Table 3. Following the SCC test, a 3-mm thick piece was sliced in parallel to the side face of the CT specimen to analyze the microstructure of the crack propagation path and the crack tip using the SEM and a JEOL-2100 transmission electron microscopy (TEM), respectively. Both sharp and blunted crack tips were selected for the TEM analysis. The TEM samples were prepared by a dual-beam focus ion beam (FIB)-SEM technique using a Quanta 200 3D system. The rest of the CT specimen was fractured by fatigue in air for SEM observation of the fracture surface as well as correction of the DCPD measurement of the crack propagation length. 3. Results 3.1 Microstructure of the fusion boundary region Delta ferrite that occurred randomly as stringers within the grains and along the grain boundaries in the HAZ of the base metal were observed by OM, shown in Fig. 3. The 308L weld metal consists of typical dendritic grains with some delta ferrite between the dendritic grain boundaries. The microhardness distribution in the FB region is shown in Fig. 4. A hardened zone with a width of approximately 5 mm was observed in the HAZ of the base metal (Fig. 4a). This resulted from the weld residual strain. The highest microhardness at about 220 HV was observed adjacent to the FB, which is about 16% higher than that of the base metal. The weld metal also shows a higher microhardness at 215 HV, which is approximately equal to that of the HAZ adjacent to the FB (Fig. 4b). The KAM and grain boundary character distribution (GBCD) in the FB region are shown in Fig. 5a and 5b, respectively. The highest residual strain corresponding to an average KAM value 5   

of 0.98 was observed adjacent to the FB in the HAZ, shown in Fig. 5a. According to the quantitative relationship between the KAM and strain for stainless steels [33], the KAM value of 0.98 corresponds to a residual strain of less than 10%. The grain boundaries shown in Fig. 5b were classified into three types: low-angle boundary (LAB) (5<<15), coincidence site lattice (CSL) grain boundary (3 ≤ Σ ≤ 29), and random high angle grain boundary (RGB). Grain boundaries in the weld metal are almost all RGBs while a high fraction of CSL grain boundaries was observed in the HAZ. Quantification of the GBCD reveals that there are about 18 % CSL grain boundaries and 75% RGBs in the HAZ, as shown in Fig. 6. Increasing the distance to the FB led to slightly increase of the CSL grain boundary but decrease of RGB fractions in the HAZ, as also shown in Fig. 6. It should be mentioned that, since cracks generally do not grow along twin boundaries, the quantification of GBCD excluded twin boundaries. 3.2 Precracking growth in the HAZ in primary water As shown in Fig. 7, a CGR of 1.7 × 10-8 mm/s was observed at the first step of the precracking under triangular waveform loading at Kmax=23 MPa√m, R = 0.5 and 0.001 Hz in primary water with a DO of less than 5 μg/L. Shifting the triangular waveform loading to 0.01 Hz and Kmax=25 MPa√m caused a higher CGR of 4.7 × 10-7 mm/s. Increasing R to 0.7 at the third step of the triangular waveform loading caused a decrease of CGR to 8.2 × 10-8 mm/s. At the fourth step, a trapezoidal waveform loading at Kmax=25 MPa√m, R=0.7 and 50s/7200s/50s was applied, which led to a very low CGR of 3.5 × 10-9 mm/s. Then the DO was increased to 200 μg/L in the last step of precracking to activate the crack growth, resulting in a higher CGR of 8.3 × 10-9 mm/s. The total crack length increment during the precracking is about 0.22 mm.  3.3 Stress corrosion crack growth in the HAZ in primary water Stress corrosion crack growth in the HAZ of 316L stainless steel in primary water with different DO, DH and K conditions are shown in Figs. 8 through 10. At the first step, a constant K at 25 MPa√m with a DO of 200 μg/L caused a low growth rate of 4.3 × 10-9 mm/s. Increasing K to 30 MPa√m in the next step resulted in a slightly higher CGR of 5.3 × 10-9 mm/s. At about 1210 h, the DO was raised to 2000 μg/L, resulting in a further increase of CGR to 9.4 × 10-9 mm/s. At 1450 h, the test was interrupted by a power outage. The test was then restarted with trapezoidal waveform loading at Kmax=31 MPa√m, R=0.7 and 50s/7200s/50s, leading to a CGR of 1.3 × 10-8 mm/s. The trapezoidal waveform loading was applied to restart the crack growth 6   

and obtain a stable CGR before switching to constant K for SCC test. The constant K at 31 MPa√m with a DO of less than 5 μg/L in the next step caused a decreased CGR of 2.9 × 10-9 mm/s over a period of approximately 250 h, as shown in Fig. 9. Switching the environment to hydrogen water chemistry with DH=2600 μg/L in the next step led to cessation of crack growth. No change in the crack growth behavior was observed after changing the DH to 860 μg/L in the next step. The crack growth was then reactivated by changing the loading mode to trapezoidal waveform, which led to a CGR of 8.0 × 10-9 mm/s. Increasing the DH to 2600 μg/L at this loading condition resulted in a slightly lower CGR of 6.5 × 10-9 mm/s. It should be mentioned that the slight decrease of the DCPD measured crack length after switching to hydrogen water chemistry is most likely associated with the dissipation of the direct current from the specimen due to the dissolution of the oxide film, which lowers the effective current through the specimen. The influent water was aerated by 200 μg/L of DO in the next step of the test, in conjunction with a change of the loading mode to constant K at 35 MPa√m. As shown in Fig. 10, a CGR of 7.5 × 10-9 mm/s was observed. The K was then ramped to 40 MPa√m over a period of 600 s, resulting in a higher CGR of 8.3 × 10−9 mm/s. The DO was switched to <5 μg/L at the last step of the test. This caused again a lower CGR of 3.1 × 10-9 mm/s. The stress corrosion crack growth rates under various K, DO and DH conditions described above are summarized in Figs. 11 and 12, respectively. 3.4 Crack propagation path and fractography SEM observation of the side face of the CT specimen after the SCC test revealed that the tip of SCC was in the HAZ with a distance of approximately 400 μm to the FB, as shown in Fig. 13a. In addition to the primary crack, a few secondary cracks were also observed. The GBCD and KAM analyses by EBSD show that the SCC is intergranular along both RGBs and CSL grain boundaries, and that there is a strain concentration adjacent to grain boundaries in the area adjacent to the SCC (Fig. 13b-c). Fracture surface of the specimen following the test as well as the relative positions on the fracture surface between the FB, the SCC tip and the precrack tip are shown in Fig. 14a-c. The tip of SCC is in HAZ and did not reach the FB (Fig. 14a), which is in consistence with the 7   

observation shown in Fig. 13a. Typical intergranular facets were observed at the SCC area (Fig. 14b-c). In addition, the average SCC propagation length is 101.5 μm, which is in agreement with the DCPD measurement shown in Figs. 8-10. 3.5 Microstructure of the crack tip Morphology and composition analysis of a blunted crack tip by TEM are shown in Fig. 15. The crack at the tip area is filled with oxide with a width of about 0.5 μm. The EDX scan across the crack plane in Fig. 15b shows that Cr-enriched oxide is formed along the crack walls next to the metal. The EDX scan along the crack plane in Fig. 15c shows that the oxide at the crack tip is enriched with Cr, while the oxide filling the crack is enriched with Ni and Fe. In addition, it is noticeable that an enrichment of Ni ahead of the crack tip in the base metal was observed. Results of the EDX scans were further confirmed by element mappings of O, Ni, Cr and Fe collected at the crack tip area, as shown in Fig. 15d. TEM observation of sharp crack tips showed that the crack had a width of about 5~10 nm (Fig. 16a). EDX scans along and across the crack plane as well as mappings of the tip area show that the oxides in the crack and at the crack tip are both enriched with Cr, but not with Fe and Ni (Fig, 16b-e). In addition, no penetrated intergranular oxidation ahead of the crack tip was observed. 4. Discussion 4.1 Stress corrosion cracking in the HAZ of 316L stainless steel The crack growth behavior described above revealed that for a stress corrosion crack propagating in the HAZ toward the FB in a stainless steel 308L-316L weld joint, the propagation rate was extremely low in normal PWR primary water chemistry. This suggests that it is unlikely the stress corrosion crack could reach the FB and penetrate into the weld metal in primary water within the 40-year lifetime of a nuclear power plant. However, adding dissolved oxygen into the water did promote the stress corrosion crack growth, leading to a potential concern for SCC in the HAZ in off-normal PWR primary water chemistry conditions. The low stress corrosion crack growth rate in the HAZ as shown in Figs 8-10 is most likely related to the relative low hardening effect as a result of the low residual strain level in the HAZ. As shown in Fig. 4a, the microhardness of the HAZ is approximately 16% higher than that of the 8   

base metal, suggesting a 16% increase of the yield strength of the HAZ. According to the relationship between the yield strength and CW level for 316L stainless steel [6, 34], the increment of the yield strength corresponds to a CW level of less than 10%. This is in consistence with the KAM analysis shown in Fig. 5. The low stress corrosion crack growth rate in the HAZ of a stainless steel weld joint with a low residual strain level in both aerated and deaerated water at high temperature was also reported in Refs [35-37]. It should be mentioned that a residual strain of more than 20% in the HAZ of a weld joint is usually expected, which causes a much higher stress corrosion crack growth rate in high temperature water [38-39]. The low residual strain level revealed in this study is most likely due to the difference in the welding process for fabrication of the weld joint. The dependence of CGR, V on K as shown in Fig 11 can be empirically described by the following relation [6]: ∝



(1)

where β is a constant describing the K dependence with a value of 0.6-3.2. For the data obtained in this study a β was found to be 1.5. This value of β is close to that observed for SCC CGRs of irradiated 316L stainless steel in high purity water at 320 C [40], but somewhat less than that for SCC CGRs of 20% CW 316L stainless steel in pure water with 2000 μg/L DO at 288 C, where β=2.3 [41]. Fig. 12 shows that the CGR increased with increasing DO in primary water, primarily due to a high corrosion rate at high electrochemical potentials caused by high DO. This is in agreement with the CGR measured from various cold worked and cast stainless steels in high temperature water [42-43]. 4.2 Crack tip microstructure and its implication to the mechanism of SCC The TEM observation and analysis of the crack tip shown in Fig. 16 indicates that an active crack tip is sharp with a width of 10-20 nm. The sharp crack is filled with Cr-enriched oxide which is protective. Further, no penetrated intergranular oxidation ahead of the crack tip in the base metal was observed. All these features suggest that the stress corrosion crack most likely propagates with a slip-oxidation mechanism [44] rather than internal oxidation. However, shear bands, a feature of the slip-oxidation mechanism, were not observed in the crack tip area. Further 9   

TEM work is needed to attain a better understanding of the SCC mechanism. The Ni enrichment ahead of the tip of a stress corrosion crack has been shown in a couple of investigations [44-49]. It was proposed that the Ni-enriched zone was attributed to selective oxidation of Cr and Fe which removed Cr and Fe preferentially from the alloy. The Ni-enriched zone could act as a barrier for oxygen diffusion and oxide growth, and thus play an important role in mitigating SCC propagation [46]. Investigations also suggested that active cracks were likely the ones without Ni buildup in the metal ahead of the crack tip [44-45]. In fact, formation of the Ni-enriched zone by a selective oxidation mechanism needs a long time [49]. With regard to an active crack, the crack tip moves forward continuously at the rate of crack growth. As such, there is insufficient time for the formation of the Ni-enriched zone ahead of the crack tip [49]. The TEM analysis shown in Fig. 15 supports these investigations. 5. Conclusion Stress corrosion crack growth in a stainless steel 308L-316L weld joint in primary water was investigated with a focus on the behavior of SCC propagation to the FB in the HAZ. The following conclusions were drawn from this investigation: (1) Stress corrosion crack growth was observed in the HAZ in off-normal PWR water chemistry conditions with DO, but not in normal PWR primary water chemistry conditions with DH. (2) The low stress corrosion crack growth rate observed is due to the low residual strain level of the HAZ, which is less than 10%. This suggests that, provided that the residual strain level in the HAZ is low, it is unlikely the SCC could reach the FB and penetrate into the weld metal in normal PWR primary water chemistry conditions within the 40-year lifetime of a nuclear power plant. (3) Stress corrosion crack growth in the HAZ of 316L stainless steel in primary water shows dependences on K and DO. (4) Observation and analysis of the crack tip by TEM suggest that the SCC in the HAZ of 316L stainless steel in primary water propagates with the slip-oxidation mechanism.

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Acknowledgement This work is jointly supported by the Hundred-Talent program of Chinese Academy of Sciences, and the National Natural Science Foundation of China (Grant No.  51571204). The authors are grateful to Dr. P.L. Andresen for his guidance in performing the CGR testing.

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[24] T.M. Angeliu, E.L. Hall, J. Sutliff, S. Sitzmen, P.L. Andresen, Strain and microstructure characterization of austenitic stainless steel weld HAZs, in: Proceedings of Corrosion 2000, March 26–31, Orlando, Florida, 2000, paper # 00186. [25] P.L. Andresen, M.M. Morra, K. Ahluwalia, SCC of Alloy 690 and its weld metals, in: Proceedings of the 15th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Colorado Springs, Colorado, August 7–11, 2011, pp. 161–176. [26] G.A. Young, N. Lewis, D.S. Morton, The stress corrosion crack growth rate of Alloy 600 heat affected zones exposed to high purity water, in: Conference on Vessel Head Penetration Inspection, Cracking, and Repairs, Gaithersburg, Maryland, March 24–26, 2003. pp. 309–324. [27] M. Ozawa, Y. Yamamoto, K. Nakata, Establishment of experimental conditions for the SCC growth rate test of Alloy 600 and Ni base weld metal in high temperature oxygenated water, 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. 639–649. [28] 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 Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Whistler, Canada, August 19–23, 2007, CD–ROM. [29] T. Arai, K. Kako, K. Watanabe, Y. Miyahara, Effect of loading direction on crack growth behavior near fusion line in low carbon stainless steel weld joints, 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. 660–670. [30] Nuclear Industry Safety Agency (NISA), Cracks on the inner surface of the welds at primary water inlet and outlet nozzle to steam generators, NISA, February 5, 2008. [31] T. Kamada, Kansai Electric’s efforts to develop aging management technologies, Presentation at ISaG, Tokyo, Japan, July 25, 2007. [32] Q.J. Peng, S. Teysseyre, P.L. Andresen, G.S. Was, Stress corrosion crack growth in type 316 stainless steel in supercritical water, Corrosion 63 (2007) 1033–1041. [33] M. Kayama, Measurement of local plastic strain distribution of stainless steel by electron 14   

backscatter diffraction, Mater. Charact. 60 (2009) 125–132. [34] K. Arioka, T. Yamada, T. Terachi, G. Chiba, Cold work and temperature dependence of stress corrosion crack growth of austenitic stainless steels in hydrogenated and oxygenated high– temperature water, Corrosion 63 (2007) 1114–1123. [35] S. Ooki, Y. Tanaka, K. Takamori, S. Suzuki, S. Tanaka, Y. Saito, T. Nakamura, T. Kato,  K. Chatani, M. Kodama, Study on SCC growth behavior of BWR core shroud, 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. 365–376. [36] S. Nouraei, D.R. Tice, K.J. Mottershead, D.M. Wright, Effects of thermo–mechanical treatments on deformation behavior and IGSCC susceptibility of stainless steels in PWR primary water chemistry, in: Proceedings of the 15th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Colorado Springs, Colorado, August 7–11, 2011, pp. 2381–2401. [37] D.R. Tice, J.W. Stairmand, H.J. Fairbrother, A. Stock, Growth testing of cold worked stainless steel in a simulated PWR primary water environment to assess susceptibility to stress corrosion cracking, in: Proceedings of the 13th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Whistler, Canada, August 19–23, 2007, CD–ROM. [38] T. Saukkonen, M. Aalto, I. Virkkunen, U. Ehrnstén, H. Hänninen, Plastic strain and residual stress distributions in an AISI 304 stainless steel BWR pipe weld, in: Proceedings of the 15th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Colorado Springs, Colorado, August 7–11, 2011, pp.2351-2367. [39] U. Ehrnstén, P. Aaltonen, P. Nenonen, H. Hänninen, C. Jansson, T. Angeliu, Intergranular cracking of AISI 316NG stainless steel in BWR environment, in: Proceedings of the 10th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Lake Tahoe, Nevada, August 5–9, 2001, CD–ROM. [40] Y. Chen, B. Alexandreanu, Y. Yang, W.J. Shack, K. Natesan, E.E. Gruber, A.S. Rao, Crack growth behavior of irradiated type 316 SS in low dissolved oxygen environment, in: Proceedings of the 15th Conference on Environmental Degradation of Materials in Nuclear Power Systems– 15   

Water Reactors, Colorado Springs, Colorado, August 7–11, 2011, pp. 1169–1182. [41] P.L. Andresen, K/Size effects on SCC in irradiated, cold worked and unirradiated stainless steel, in: Proceedings of the 11th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Stevenson, Washington, August 10–14, 2003, pp. 870– 886. [42] M.B. Toloczko, P.L. Andresen, S.M. Bruemmer SCC crack growth of cold–worked type 316SS in simulated BWR oxidizing and hydrogen water chemistry conditions, in: Proceedings of the 13th Conference on Environmental Degradation of Materials in Nuclear Power Systems– Water Reactors, Whistler, Canada, August 19–23, 2007, CD–ROM. [43] T. Yamada, T. Terachi, T. Miyamoto, K. Arioka, Crack growth behavior of welded and cast stainless steels in hydrogenated and oxygenated high–temperature water, 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. 684–689. [44] S.M. Bruemmer, L.E. Thomas, Comparison of IGSCC crack–tip characteristics produced in BWR oxidizing water and BWR hydrogen water chemistry conditions, in: Proceedings of the 13th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Whistler, Canada, August 19–23, 2007, CD–ROM. [45] L.E. Thomas, D. Edwards, K. Asano, S. Ooki, S.M. Bruemmer, Crack–tip characteristics in BWR service components, in: Proceedings of the 13th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Whistler, Canada, August 19–23, 2007, CD–ROM. [46] K. Kruska, S. Lozano-Perez, D.W. Saxey, T. Terachi, T. Yamada, G.D.W. Smith, Nanoscale characterization of grain boundary oxidation in cold–worked stainless steels, Corros. Sci. 63 (2012) 225–233. [47] N.K. Das, K. Suzuki, K. Ogawa, T. Shoji, Early stage SCC initiation analysis of FCC Fe– Cr–Ni ternary alloy at 288 °C: a quantum chemical molecular dynamics approach, Corros. Sci. 51 (2009) 908–913. [48] T. Terachi, K. Fujii, K. Arioka, Microstructural characterization of SCC crack tip and oxide film for SUS316 stainless steel in simulated PWR water at 320 °C, J. Nucl. Sci. Technol. 42 16   

(2005) 225–232. [49] E.P. Simonen, L.E. Thomas, S.M. Bruemmer, Defining the role of solute diffusion kinetics at crack tips in stainless steel during irradiation in light–water reactors, in: Proceedings of the 11th Conference on Environmental Degradation of Materials in Nuclear Power Systems–Water Reactors, Stevenson, Washington, August 10–14, 2003, pp. 1062–1070.

Figure captions Figure 1. Schematic drawing showing how the CT and microstructure-analysis specimens were extracted from the weld joint. Figure 2. Location of the tips of the precrack on side faces of the CT specimen before the test in primary water. Figure 3. OM observation of the microstructure of the fusion boundary region in the stainless steel 308L-316L weld joint. Figure 4. Microhardness distribution in the stainless steel 308L-316L weld joint. (a), the overall distribution, and (b), the distribution in the fusion boundary region. Figure 5. KAM (a) and grain boundary character (b) distributions in the fusion boundary region of the stainless steel 308L-316L weld joint. In Fig. b, the black lines represent random high angle grain boundaries, the red represent CSL grain boundaries and the white represent the low angle boundaries. Figure 6. Grain boundary character distribution versus distance to the fusion boundary in the HAZ of 316L stainless steel. Figure 7. Crack growth in the HAZ of 316L stainless steel during in-situ precracking in primary water.

17   

Figure 8. Stress corrosion crack growth in the HAZ of 316L stainless steel in primary water in different K and DO conditions. Figure 9. Stress corrosion crack growth in the HAZ of 316L stainless steel in primary water in various DH conditions. Figure 10. Stress corrosion crack growth in the HAZ of 316L stainless steel in various DO and K conditions. Figure 11. Effect of K on the stress corrosion crack growth in the HAZ of 316L stainless steel in primary water with 200 μg/L of DO at 320 °C. Figure 12. Effect of water chemistry on stress corrosion crack growth in the HAZ of 316L stainless steel in primary water at 320 °C. Figure 13. SEM observation and EBSD analysis of the crack propagation path after the SCC test. (a), crack morphology on the side face of the specimen, (b), grain boundary character distribution, and (c), KAM distribution. Crack propagation is from bottom to top. Figure 14. SEM observation of the fracture surface of the specimen after the SCC test. (a), outline of the fracture surface showing the relative positions between the FB, the SCC tip and the precrack tip, (b) and (c), observation of the SCC region at two magnifications. Crack propagation is from bottom to top. Figure 15. Observation and analysis of a blunted crack tip by TEM. (a), crack tip morphology, (b), composition profile across the crack plane, (c), composition profile along the crack plane, and (d), element mappings for the crack tip area. Figure 16. Observation and analysis of a sharp crack tip by TEM. (a), Crack tip morphology, (b), composition profile along the crack plane, (c) and (d), composition profiles across the crack plane, and (e), element mappings for the crack tip area. 18   

Figure 1. Schematic drawing showing how the CT and microstructure-analysis specimens were extracted from the weld joint.

19   

Figure 2. Location of the tips of the precrack on side faces of the CT specimen before the test in primary water.

20   

Figure 3. OM observation of the microstructure of the fusion boundary region in the stainless steel 308L-316L weld joint.

21   

 

FB

FB 260

(a) 316L

308L

HAZ

220

200

180

180 -10

-5

0

5

Distance to the FB (mm)

Measurement 1 Measurement 2 Measurement 3 Measurement 4

220

200

-15

(b)

240

HV0.1

240

HV0.1

260

316L

-2

10

-1

308L

0

1

Distance to the FB (mm)

2

Figure 4. Microhardness distribution in the stainless steel 308L-316L weld joint. (a), the overall distribution, and (b), the distribution in the fusion boundary region.

22   

Figure 5. KAM (a) and grain boundary character (b) distributions in the fusion boundary region of the stainless steel 308L-316L weld joint. In Fig. b, the black lines represent random high angle grain boundaries, the red represent CSL grain boundaries and the white represent the low angle boundaries.

23   

Grain boundary percentage (%)

100

LAB CSL grain boundary RGB

90 80 70 60 50 40 30 20 10 0

0.4

1.2

2.0

Distance to the FB (mm)

2.8

Figure 6. Grain boundary character distribution versus distance to the fusion boundary in the HAZ of 316L stainless steel.

24   

Crack length (mm) 13.00

 

12.95

12.90

12.85

12.80

12.75

13.05

12.70

12.65 0

8.210-8mm/s

4.710-7mm/s

-8

1.710 mm/s

200

Time (h) 400

water.

25 

3.510-9mm/s

600

To, Trap. loading, 50s/7200s/50s Kmax=25 MPam R=0.7, DO=200 μg/L

To, Trap. loading, 50s/7200s/50s Kmax=25 MPam R=0.7, DO<5 μg/L

To, R=0.7 f=0.01Hz Kmax=25 MPam DO<5 μg/L

Start Tri. loading at R=0.5, f=0.001Hz, Kmax=23 MPam, DO<5 μg/L To R=0.5 f=0.01Hz Kmax=25 MPam, DO<5 μg/L

 

8.310-9mm/s

Primary water, 320 oC 800

Figure 7. Crack growth in the HAZ of 316L stainless steel during in-situ precracking in primary

12.99

12.95 12.94 12.93

1.310-8mm/s

12.92 12.91

Constant K at 31MPam DO=2000 μg/L

12.96

Constant K at 30MPam DO=200 μg/L

12.97

Constant K at 25MPam DO=200 μg/L

Crack length (mm)

12.98

Restart test, trap. loading R=0.7 50s/7200s/50s Kmax=31MPam DO=2000 μg/L

 

5.310 mm/s -9

4.310-9mm/s

9.410-9mm/s

Primary water, 320oC

12.90 800

1000

1200

1400

1600

Time (h) Figure 8. Stress corrosion crack growth in the HAZ of 316L stainless steel in primary water in different K and DO conditions.

26   

 

2.910-9mm/s

12.94

No crack growth

To Trap. loading R=0.7 50s/7200s/50s Kmax=31MPam DH=2600 μg/L

To Trap. loading R=0.7 50s/7200s/50s Kmax=31MPam DH=860 μg/L

12.95

Constant K at 31 MPam DH=860 μg/L

12.96

Constant K at 31MPam DO<5 μg/L

Crack length (mm)

12.97

Constant K at 31MPam DH=2600 μg/L

12.98

6.510-9mm/s 8.010-9mm/s

Primary water, 320oC

12.93

1600

1800

2000 2200 2400 2600 Time (h) Figure 9. Stress corrosion crack growth in the HAZ of 316L stainless steel in primary water in various DH conditions.

27   

12.97 12.96

Constant K at 40 MPam DO<5 μg/L

12.98

Constant K at 40 MPam DO=200 μg/L

Crack length (mm)

12.99

Constant K at 35 MPam DO=200 μg/L

 

8.310 mm/s -9

12.95

7.510-9 mm/s

3.110-9 mm/s

Primary water, 320 oC

12.94

2600 2700 2800 2900 3000 3100 3200 3300

Time (h)

Figure 10. Stress corrosion crack growth in the HAZ of 316L stainless steel in various DO and K conditions.

28   

Crack growth rate (pm/s)

50

o

primary water, 320 C, DO = 200 μg/L

1.5

10

1

V  K

20

25

30

35

K (MPam)

40

45

Figure 11. Effect of K on the stress corrosion crack growth in the HAZ of 316L stainless steel in primary water with 200 μg/L of DO at 320 °C.

29   

 

Crack growth rate (pm/s)

15

10

DH= 2600 or 860 μg/L, constant K at 31 MPam <5 μg/L DO, constant K at 31M Pam 200 μg/L DO, constant K at 30 MPam 2000 μg/L DO, constant K at 31 MPam

primary water, 320 oC

5

No crack growth

0 DH

<5 μg/L DO 200 μg/L DO 2000 μg/L DO

Figure 12. Effect of water chemistry on stress corrosion crack growth in the HAZ of 316L stainless steel in primary water at 320 °C.

30   

Figure 13. SEM observation and EBSD analysis of the crack propagation path after the SCC test. (a), crack morphology on the side face of the specimen, (b), grain boundary character distribution, and (c), KAM distribution. Crack propagation is from bottom to top.

31   

Figure 14. SEM observation of the fracture surface of the specimen after the SCC test. (a), outline of the fracture surface showing the relative positions between the FB, the SCC tip and the precrack tip, (b) and (c), observation of the SCC region at two magnifications. Crack propagation is from bottom to top.

32   

Figure 15. Observation and analysis of a blunted crack tip by TEM. (a), crack tip morphology, (b), composition profile across the crack plane, (c), composition profile along the crack plane, and (d), element mappings for the crack tip area.

33   

Figure 16. Observation and analysis of a sharp crack tip by TEM. (a), Crack tip morphology, (b), composition profile along the crack plane, (c) and (d), composition profiles across the crack plane, and (e), element mappings for the crack tip area.

34   

Table 1 Chemical composition (wt%) of the base and weld metals of the weld joint C

Si

Mn

Mo

Ni

Cr

Fe

316L

0.025

0.52

1.71

2.4

11.7

17.9

Bal.

308L

<0.3

0.35

0.9

-

11

20

Bal.

Table 2 Loading and environment conditions for the in-situ precracking steps in primary water Steps #1 #2 #3 #4 #5

Loading mode Triangular waveform, R = 0.5, 0.001 Hz, Kmax = 23 MPa√m Triangular waveform, R = 0.5, 0.01 Hz, Kmax = 25 MP√m Triangular waveform, R = 0.7, 0.01 Hz, Kmax = 25 MPa√m Trapezoidal waveform, R = 0.7, 50s/7200s/50s, Kmax = 25 MPa√m Trapezoidal waveform, R = 0.7, 50s/7200s/50s, Kmax = 25 MPa√m

35   

Environment 320 °C, DO < 5 μg/L 320 °C, DO < 5 μg/L 320 °C, DO < 5 μg/L 320 °C, DO < 5 μg/L 320 °C, DO = 200 μg/L

Table 3 Loading and environment conditions for each step of the SCC test Steps

Loading mode

Environment

#1

Constant K at 25MPa√m

320 °C, DO = 200 μg/L

#2

Constant K at 30MPa√m

320 °C, DO = 200 μg/L

#3

Constant K at 31MPa√m

320 °C, DO = 2000 μg/L

#4

Trapezoidal waveform, R = 0.7, 50s/7200s/50s, Kmax=31MPa√m

320 °C, DO = 2000 μg/L

#5

Constant K at 31MPa√m

320 °C, DO < 5 μg/L

#6

Constant K at 31MPa√m

320 °C, DH = 2600 μg/L

#7

Constant K at 31MPa√m

320 °C, DH = 860 μg/L

#8 #9

Trapezoidal waveform, R = 0.7, 50s/7200s/50s, Kmax = 31MPa√m Trapezoidal waveform, R =0.7, 50s/7200s/50s, Kmax = 31MPa√m

320 °C, DH = 2600 μg/L

#10

Constant K at 35 MPa√m

320 °C, DO = 200 μg/L

#11

Constant K at 40 MPa√m

320 °C, DO = 200 μg/L

#12

Constant K at 40 MPa√m

320 °C, DO < 5 μg/L

36   

320 °C, DH = 860 μg/L