Characterization of microstructure, local deformation and microchemistry in Alloy 690 heat-affected zone and stress corrosion cracking in high temperature water

Journal of Nuclear Materials 465 (2015) 471e481

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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat

Characterization of microstructure, local deformation and microchemistry in Alloy 690 heat-affected zone and stress corrosion cracking in high temperature water Zhanpeng Lu a, b, *, Junjie Chen a, Tetsuo Shoji c, Yoichi Takeda c, Seiya Yamazaki c a b c

Institute of Materials Science, School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China State Key Laboratory of Advanced Special Steels, Shanghai University, Shanghai 200072, China New Industry Creation Hatchery Center, Tohoku University, Sendai 980-8579, Japan

a r t i c l e i n f o

a b s t r a c t

Article history: Received 13 April 2015 Received in revised form 3 June 2015 Accepted 12 June 2015 Available online 18 June 2015

With increasing the distance from the weld fusion line in an Alloy 690 heat-affected zone, microhardness decreases, kernel average misorientation decreases and the fraction of S3 boundaries increases. Chromium depletion at grain boundaries in the Alloy 690 heat-affected zone is less significant than that in an Alloy 600 heat-affected zone. Alloy 690 heat-affected zone exhibits much higher IGSCC resistance than Alloy 600 heat-affected zone in simulated pressurized water reactor primary water. Heavily cold worked Alloy 690 exhibits localized intergranular stress corrosion cracking. The effects of metallurgical and mechanical properties on stress corrosion cracking in Alloy 690 are discussed. © 2015 Elsevier B.V. All rights reserved.

Keywords: Nickel-base alloy Stress corrosion cracking Welding Heat-affected zone Pressurized water reactor High temperature water

1. Introduction Ni-based alloys and weld metals are used extensively for fabricating structural components in pressurized water reactors (PWRs). Stress corrosion cracking (SCC) in Ni-base Alloy 600 and its weld metals is one of the main issues of material degradation in PWR coolants [1e7]. Lots of crack growth rate (CGR) data for Alloy 600 base metals and weld metals in the as-received conditions or in the cold-worked conditions have been measured in simulated PWR primary water environments [1e6]. Many properties in a weld heataffected zone (HAZ) are different from those in the base metal and can contribute to the change of SCC resistance. The thermal cycles from welding process can result in residual stress and strain and the changes of microstructure, carbide distribution and local material chemistry and therefore affect the primary water stress corrosion cracking (PWSCC) behavior [8e10]. The inhomogeneity in metallurgical and mechanical properties in weld HAZs is schematically

* Corresponding author. Institute of Materials Science, School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China. E-mail address: [email protected] (Z. Lu). http://dx.doi.org/10.1016/j.jnucmat.2015.06.025 0022-3115/© 2015 Elsevier B.V. All rights reserved.

shown in Fig. 1，which is affected by welding methods, welding procedures and parameters. Young et al. [8] have investigated the SCC growth behavior of an Alloy 600 weld HAZ in high temperature pure water and the effects of microstructure, test temperature and electrochemical potential as function of dissolved hydrogen concentration. They reported that the CGRs in the Alloy 600 HAZ were about 30 times faster than those in the Alloy 600 base metal tested under the same conditions [8]. Lu et al. [9,10] have measured the grain boundary character distribution, grain boundary carbide distribution, grain boundary chemistry and local deformation in an Alloy 600 HAZ and correlated to its SCC behavior in simulated PWR primary water. The CGR in the Alloy 600 HAZ can be more than 20 times higher than that in its base metal. The effect of strainhardening in weld HAZs can be simulated partly by using cold worked Alloys. The grain boundary character distribution and local deformation in Alloy 600 cold worked to various degrees were measured. With increasing the degree of cold work, the hardness P and kernel average misorientation increase while the fraction of 3 boundaries decreases. The PWSCC behavior in the Alloy 600 HAZ and cold worked Alloy 600 were investigated and compared [9,10]. Alloy 690 and its weld metals such as Alloys 52, 52M, 152 have been used in new PWR plants or used to replace previously used

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Fig. 1. Schematic of a weld exhibiting metallurgical and mechanical characteristics in HAZs.

Alloy 600 and its weld metals such as Alloys 82 and 182. It is important to evaluate the SCC resistance of Alloy 690 and its weld HAZ in PWR primary water environments and related influencing factors on SCC. The complexity related to bulk chemistry, grain boundary chemistry, grain boundary types and ratios, mechanical properties and other factors in Ni-base weld HAZs such as Alloy 600 HAZ and Alloy 690 HAZ has been pointed out [9e13]. Straining hardening has been identified to be one of the factors contributing to SCC in stainless steel and Ni-base alloy weld HAZs [9e16]. Andresen [12] reported the SCC growth rates under various loading modes for an Alloy 690/EN52M narrow gap weld HAZ in a simulated PWR primary water at 360
C. The CGR in the HAZ was 7  1012 m/s under the constant loading. The SCC growth was significantly enhanced by incorporating an unloading-reloading step in the trapezoidal wave loading mode, and the enhancement factor decreased with increasing the holding time during the trapezoidal wave loading. Alexandrean et al. [13] measured the CGRs under cyclic and constant loading modes for Alloy 690 HAZ in a PWR environment at 320
C, and reported the measured CGRs to be in the range of 3.36  1012 to 7.74  1012 m/s under constant loading at K of about 30 MPa m0.5. It was found that different HAZs could exhibit different CGRs, considering the above reported CGR data and the effect of temperature on SCC growth. Different microstructures, bulk chemistries and microchemistries and local mechanical properties in weld HAZs as the results of different welding procedures and parameters could contribute to the observed differences in SCC growth behavior of HAZs. The effects of cold work and the details of cold work such as rolling methods and orientations on SCC of Alloy 690 were investigated. Cold work could significantly accelerate the PWSCC growth rate of Alloy 690 [12e16]. Scatted CGR data for asreceived Alloy 690 and cold worked Alloy 690 from different heats were observed. In the present work, the hardness distribution, microstructure and microchemistry at grain boundaries of an Alloy 690 weld HAZ are characterized. PWSCC growth behaviors of Alloy 690 HAZ and cold worked Alloy 690 in PWR primary water environments are investigated and compared to the results of Alloy 600 HAZ and cold worked Alloy 600. The effects of strain-hardening, microstructure and the bulk/grain boundary chemistry on SCC of Ni-base alloy HAZs are discussed. 2. Experimental details 2.1. Test materials Alloy 690 base metal and weld heat-affected zone in an Alloy

Table 1 Chemical compositions of Alloy 690 (wt. %). C

Si

Mn

S

P

Cu

Ni

Cr

Fe

0.02

0.11

0.16

0.001

0.006

0.01

59.42

29.35

10.13

690/Alloy 52 weld were used in the experiments. The weld was prepared by Shielded Metal Arc Welding (SMAW) using thermally treated Alloy 690 base metal (called Alloy 690TT) and Alloy 52 weld metal. The Alloy 690 was obtained by mill annealing (MA) at 1100
C for 1 h followed by air cooling. Thermal treatment (TT) was performed at 700
C for 15 h. Alloy 690TT was used for preparing the weld block. The chemical compositions and mechanical properties of Alloy 690TT base metal are shown in Tables 1 and 2, respectively. The schematic of the Alloy 690TT weld block is shown in Fig. 2. Fusion line A1 was produced by the first welding. Fusion lines A and B were produced by the second welding after removal a part of the weld metal of the first welding. If not specially mentioned, the weld fusion line refers to fusion line A throughout the text. A 40% cross-rolled Alloy 690TT block, called as 40% CR Alloy 690, was prepared and used in the SCC tests for understanding the effect of strain-hardening on SCC. The CR procedure was similar to that has been used for Alloy 600 [9,10]. The parameters for the CR are listed in Table 3. 2.2. Characterization of microstructure and local deformation The values of Vickers hardness (HV) in the Alloy 690TT HAZ and in the 40% Cr Alloy 690TT were measured at a load of 9.8 N and a holding time of 15 s. The HV in the Alloy 690TT HAZ was measured as a function of the distance from the weld fusion line A. Microstructure and grain boundary properties were characterized by electron back scattering diffraction (EBSD). Local deformation was evaluated by HV and kernel average misorientation (KAM) measurement by EBSD technique 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

Table 2 Mechanical properties of Alloy 690 base metal at RT. Yield strength, MPa

Tensile strength, MPa

Elongation ratio, %

256

619

59.9

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in a high vacuum connected to the AES chamber [10]. The specimen size and the location of SSRT specimen are shown in Fig. 3. Hydrogen charging was performed in 36
C 0.5 mol/L sulfuric acid containing 1 g/L thiourea with nitrogen gas purging at a constant current density of 5 mA/cm2. The hydrogen charging time was 1269.3 h. The strain rate for SSRT was 6.4  107 s1. The gauge part of the SSRT specimen located at about 2 mm away from the fusion line A in the Alloy 690TT HAZ. The hydrogen-charged specimen was electroplated with copper to prevent the outflow of hydrogen from the specimen. After electroplating, the sample was immediately introduced into the high vacuum chamber with a SSRT machine, so that IG surface analysis could be performed without contamination. The chemical compositions of various locations on the IG facets of the fracture surface were analyzed by AES. AES beam parameters for analysis were 10 kV and 10 nA and the beam size was around 40e50 nm. Fig. 2. The schematic of an Alloy 690TT/Alloy 52 weld used for the analysis of material properties and for taking the sample for SCC tests in a simulated PWR primary water environment. Fusion line A1 was produced by the first welding. Fusion lines A and B were produced by the second welding after removal a part of the weld metal of the first welding.

was analyzed using OIM-Analysis software provided by TSL, Co. Ltd [17]. The SEM/EBSD technique provides KAM as a parameter to evaluate the plastic strain for a given point. The KAM is defined for a given point as the numerical average misorientation of that point with all of its six neighbors [17]. Acceleration voltage of SEM beam for the EBSD measurement was 25 kV. The surface was finished by polishing with 1 mm diamond paste followed by electropolishing using HClO4 þ C2H5OH electrolyte in order to get relatively smooth surface free from surface hardening caused by the mechanical polishing [10].

2.3. Characterization of grain boundary chemistry Slow strain rate test (SSRT)/Auger Electron Spectroscopy (AES) technique was used to measure the grain boundary chemistry in Alloy 690TT HAZ. Investigation of grain boundary chemistry was carried out by analyzing intergranular (IG) fracture surfaces that were formed by SSRT of a hydrogen-charged miniatured specimen

2.4. Crack growth rate tests in a simulated PWR primary water environment Contoured double cantilever beam (CDCB) specimens were fabricated from an Alloy 690 HAZ block and a 40% CR Alloy 690TT block. The design and location of the Alloy 690 HAZ CDCB specimen (25 mm thick) for SCC growth test was similar to that for the Alloy 600 HAZ CDCB specimen [9，10]. According to the HV results, the notch tip of the crack growth test specimen was designed to be in the heavily hardened area in the Alloy 690 HAZ. The notch direction was parallel to the fusion line A. The notch tip was about 2 mm (x ¼ 2 mm) away from the fusion line A in the Alloy 690 HAZ specimen. The notch of the 40% CR Alloy 690 CDCB specimen (12.5 mm thick) was prepared in the transverse-longitudinal (T-L) orientation, which was similar to that described in a previous paper for 40% CR Alloy 600 [9]. CGR tests were performed in a simulated PWR primary water (B: 1200 ppm (wt.) as H3BO3, Li: 2.0 ppm (wt.) as LiOH, DO < 5 ppb, DH ¼ 30 cm3 (STP) H2/kg H2O) at 340
C. The inlet conductivity was about 20 mS/cm at RT. The specimens were first air-fatigued with a Kmax < 20 MPa m0.5 and load ratio (R ¼ Kmin/Kmax) R ¼ 0.2, then grooved 5% on each side. After in-situ pre-cracking in the simulated PWR primary water, trapezoidal (Trap.) loading or constant loading (CL) was applied during the SCC

Table 3 Procedures for preparing 40% CR Alloy 690. Material

Percent of reduction in thickness in each pass, x

Number of rolling passes, n

Total reduction in thickness, x*n

Alloy 690TT

5%

8

40%

Fig. 3. Schematic illustrations of SSRT-AES specimen (a) specimen design and (b) location of the SSRT specimen in Alloy 690TT HAZ. Unit is in mm.

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direction. HV values in the 40% CR Alloy 690 were much higher than that in the Alloy 690 base metal and in the Alloy 690 HAZ. 3.2. Grain boundary character distribution

Fig. 4. The loading patterns for the precracking in high temperature water followed by the SCC test in high temperature water.

tests, as schematically shown Fig. 4. Details of loading modes and test periods after in situ precracking under triangular wave loading are shown in Table 4. 3. Results and discussion 3.1. Hardness distribution The distribution of HV as well as the location of CDCB specimen for the CGR test in the Alloy 690 HAZ is shown in Fig. 5. The heavily hardened area in the Alloy 690 HAZ indicated by hardness measurements was within a distance of about 2 mm away from the fusion line A, in which the HV was around 220e250. The HV in the Alloy 690 HAZ generally decreased with increasing the distance (if more than about 2.0 mm) from the fusion line A. The distribution of HV in the 40% CR Alloy 690 is shown in Fig. 6. Values of HV across the thickness direction were in the range of 330e360, showing that the cross-rolled plate was relatively uniform in the thickness

The inverse pole figures by EBSD measurements for the Alloy 690 HAZ near the weld fusion line are shown in Fig. 7. The EBSD data are lightly cleaned up to show more clearly of grain boundaries. Large size dendrite microstructure was observed in the weld metal Alloy 52. In the HAZ within a distance of about 0.3e0.4 mm from the fusion line, grain size was significantly larger than that in the base metal. If the distance from the fusion line was more than 1 mm, the grain size was close to that in the base metal, implying that grain size change by welding process occurred only in a narrow region of the Alloy 690 HAZ. The region of notch tip of Alloy 690 HAZ CDCB specimen was about 2 mm in the X-direction away from the fusion line A, in which the grain size was similar to that in the base metal. More detailed EBSD results of Alloy 690TT HAZ as function of the distance from the weld fusion line A are shown in Fig. 8. Fractions of S3 boundaries vs. distance from the weld fusion line A in the Alloy 690 HAZ are shown in Fig. 9. The fractions at specific KAM values vs. distance from the weld fusion line A in the Alloy 690 HAZ are shown in Fig. 10. Fractions of S3 boundaries increased and KAM decreased with increasing the distance from the weld fusion line. 3.3. Grain boundary microchemistry The SEM photo and the illustration for the fracture surface of an Alloy 690 HAZ specimen after hydrogen charging in a sulfuric acid solution with thiourea and then SSRT to fracture in a high vacuum chamber are shown in Fig. 11. Clear intergranular facets were produced by such a procedure, which were subsequently analyzed by

Table 4 SCC test procedures in a simulated PWR primary water at 340
C. Material and specimen

Loading level, loading mode and time

Alloy 690 HAZ CDCB specimen

(1) CL: 628.8 h, (2) Trap. Loading (R ¼ 0.7, 60s unload/reload, 72 h holding time): 3730.7 h. The total SCC test time was 4359.5 h. K and Kmax were 30 MPa m0.5. (1) CL: 118 h, (2) Trap. Loading (R ¼ 0.7, 60s unload/reload, 72h holding): 3932 h. The total SCC time was 4050 h. K and Kmax were 30 MPa m0.5.

40% CR Alloy 690 CDCB specimen

Fig. 5. The distribution of HV as well as the alignment of the specimen for the CGR test in the Alloy 690TT HAZ in an Alloy 690TT/Alloy 52 weld used in the experiments.

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Fig. 6. The distribution of HV in the 40% CR Alloy 690TT along the thickness direction at two locations, lines 1 and 2.

AES to quantify the distribution of different elements at grain boundaries. The ratio of the intergranular area on the whole fracture surface was dependent on both material properties and the hydrogen charging condition/time. A high resolution SEM photograph and AES results of typical intergranular facets on the fracture surface are shown in Fig. 12. Carbides, phosphorous, sulfur and boron were found on the intergranular facets of Alloy 690 HAZ. Crdepletion at grain boundaries was not significant in the Alloy 690 HAZ, as shown in Fig. 12 (b)e(d). Cr-depletion at grain boundaries has been found to be more significant in the Alloy 600MA HAZ than that in the Alloy 690 HAZ [10]. 3.4. SCC growth behavior SEM morphologies of the fracture surfaces for Alloy 690 HAZ

475

after CGR testing in the simulated PWR primary water at 340
C for 4359.5 h are shown in Fig. 13. The SCC fracture surface of Alloy 600MA HAZ under the same test condition for the same test period [10] are shown in Fig. 14 for comparison. Very tiny cracking on the fracture surface of the Alloy 690 HAZ specimen was observed, showing rather high SCC resistance of the Alloy 690 HAZ specimen. Significantly long intergranular SCC cracks were observed on the fracture surface of the Alloy 600 HAZ, showing a much higher CGR in the Alloy 600 HAZ than that in the Alloy 690 HAZ. SEM morphologies of the fracture surface of 40% cross-rolled Alloy 690 specimen after SCC testing in a simulated PWR primary water at 340
C for 4050 h are shown in Fig. 15, showing mixed intergranualr-transgranular cracking mode with local intergranualar cracks. Cold work by cross-rolling significantly increased the SCC sensitivity of Alloy 690. It has been reported that the SCC growth rate in 40% CR Alloy 600 in a simulated PWR primary water at 320
C was about 1.82  1010 m/s. After calibration of CGR with apparent activation energy of 130 kJ/mol [1], the PWSCC CGR in 40% CR Alloy 600 was much higher than that in 40% CR Alloy 690 at the same test temperature.

3.5. Comparison of SCC behavior of Alloy 690 HAZ and Alloy 600 HAZ Since the strain-hardening and the resultant changes of hardP ness, fraction of 3 boundaries, and KAM distribution in Alloy 690 HAZ and Alloy 600 HAZ are generally similar [9e11], it is thought that the bulk or grain boundary chemical composition would contribute significantly to the different SCC behaviors of Alloy 690 HAZ and Alloy 600 HAZ in simulated PWR environments. Similarly, difference in chemical composition could contribute to the significantly higher PWSCC resistance in cold worked Alloy 690TT than that in cold worked Alloy 600. The chemical composition at grain boundaries and the resultant grain boundary corrosion or oxidation kinetics would play an important role in intergranular SCC. The differences of SCC growth behavior of Alloy 600 HAZ and Alloy 690 HAZ are analyzed in terms of the oxidation kinetics as the results of bulk and grain boundary chemical compositions. SCC growth rates (da/dt) of austenitic alloys in high temperature water can be correlated with the oxidation charge density (Qd) occurring during the crack tip film degradation period (td), according to Faraday's law [18e22].

da ¼ dt



   Ma Q $ d rzF td

(1)

where Ma is the atomic weight, r is the density of the material, z is the change of valence due to oxidation, and F is Faraday's constant. The physical effect of strain on film integrity is expressed by  crack tip strain rate, εct [18e22].

. td ¼ εd εct

Fig. 7. Inverse pole figure (IPF) results of Alloy 690TT HAZ adjacent to the weld fusion line obtained by EBSD measurements (step size ¼ 6 mm).

(2)

where td is the time period for the onset of film degradation, εd is  the film degradation strain, and εct is the crack tip strain rate. If the oxidation rate is dominantly controlled by transient oxidation process such as repassivation process or metal oxidation process, the instantaneous oxidation rate is time-dependent, even though the overall CGR for a long time span can be regarded to be constant. A general transient oxidation kinetics equation is proposed based on a power law oxidation rate decay process that is generally applicable to both repassivation kinetics and quasi-solid state oxidation kinetics [23]. Then the crack growth rate equation can be written as,

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Fig. 8. EBSD results of Alloy 690TT HAZ adjacent to the weld fusion line (step size ¼ 2 mm).

  m da ¼ ka $ εct dt

(3)

where a is the crack length, da/dt is the crack growth rate, ka is crack tip oxidation rate constant, m is the slope of the oxidation rate decay curve. The definition of ka depends on the rate-determining step for the crack tip oxidation process. Several SCC mechanisms have been proposed concerning the element processes and rate-controlling laws, such as slipdissolution/oxidation, internal oxidation, quasi-solid state

oxidation mechanism [1,20e24]. Based on the slip-dissolution/ oxidation mechanism [20,21], there is



Ma $i0 ka ¼ z$r$F$ð1  mÞ



t0 εf

!m (4)

where i0 is the active surface oxidation current density, t0 is the time for the onset of current decay, m is the slope of the current decay curve, and εf is the film rupture strain. For quasi-solid state oxidation mechanism [23,24], there are

Z. Lu et al. / Journal of Nuclear Materials 465 (2015) 471e481

i h ka ¼ ðk1 Þð1mÞ $ðεd ÞðmÞ in  om da h ¼ ðK1 Þð1mÞ $ðεd ÞðmÞ εct dt

477

(5)

(6)

The quasi-solid state oxidation kinetics law and related constants, k1, m1 and m, are represented by the following equations [23e25].

Lm1 ¼ k1 t  m¼

Fig. 9. Fractions of S3 boundaries vs. distance from the weld fusion line A in the Alloy 690TT HAZ.

1

(7) 1 m1

 (8)

where L is the film thickness, and k1 is the oxidation rate constant, and t is time. There are m1 ¼ 2 and m ¼ 0.5 for the parabolic oxidation kinetics, which follow Wagner's theory of oxidation [25],

L 2 ¼ k1 t

(9)

Combining the crack tip strain-gradient theory and different crack tip asymptotic fields, theoretical crack growth rate equations are derived and applied to modeling of SCC growth rates of austenitic alloys in high temperature water environments [23,24]. Intergranular SCC of alloys in high temperature water is closely related to the local oxidation kinetics at grain boundaries. Intergranular SCC growth rate can be expressed by Eq. (10),

ð1mÞ  ðmÞ in  om da h ¼ k1;GB εct;GB $ εd;GB dt

Fig. 10. KAM vs. distance from the weld fusion line A in the Alloy 690TT HAZ. (step size ¼ 2 mm).

(10)

where k1,GB is the oxidation rate constant at grain boundaries, εd;GB  is the film degradation strain at grain boundaries, and εct;GB is the crack tip strain rate at grain boundaries. A significant effect of Cr content in Ni-base alloys on the SCC behavior in simulated PWR environments has been reported by Morton et al. [26]. Variations in chemical compositions at grain boundaries such as enrichment/depletion of specific elements are schematically shown in Fig. 16, which would affect the local oxidation kinetics in terms of k1,GB, film roughness in terms of εd;GB  and εct;GB . In Fig. 16, element A depletes and element B enriches at grain boundaries. The oxidation rate constants for Ni-base alloys in high

Fig. 11. The SEM morphology a) and the illustration b) of the fracture surface for the hydrogen charged (for 1269.3h) Alloy 690TT HAZ SSRT specimen.

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Fig. 12. SEM photograph a) and AES spectra b), c) and d) of a typical inter-granular facet on the fracture surface of the Alloy 690TT HAZ after hydrogen-charging in solution and then SSRT to fracture in vacuum.

Fig. 13. SEM morphologies of the fracture surface for Alloy 690TT HAZ after SCC testing in a simulated PWR primary water at 340
C for 4359.5 h a), b) and c) are different locations on the fracture surface.

Z. Lu et al. / Journal of Nuclear Materials 465 (2015) 471e481

479

Fig. 16. Schematic of depletion of element A or enrichment of element B at a grain boundary. Fig. 14. SEM photos and the illustration of the fracture surface of Alloy 600MA HAZ specimen after testing in simulated PWR primary water at 340
C for 4359.5 h.

temperature water environments are functions of alloy compositions, temperature, water chemistry and many other factors. Oxidation behaviors of Alloys 600, 182 and 690 in simulated PWR

primary water and of Alloy 600 in hydrogenated steam have been investigated [27e32]. The values of k1 calculated with Eq. (9) based on the available immersion data from Refs. [27e32] are shown in Fig. 17. Results show that the values of k1 for Alloy 600

Fig. 15. SEM morphologies of the fracture surface for 40% cross-rolled Alloy 690 TT after SCC testing in a simulated PWR primary water at 340
C for 4050 h a) whole surface, and b) and c) for typical SCC regions.

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Fig. 17. Calculated parabolic rate constant k1 for nickel-base alloys in simulated PWR primary water and in steam at different temperatures [27e32].

and 182 in simulated PWR primary water and steam increase with increasing temperature, following a thermal-activation mechanism. The values of k1 for Alloys 600 and 182 with Cr contents of about 15% (wt.) in simulated PWR primary water environments are significantly higher than that for Alloy 690 with Cr contents of about 30% [27e33]. The dependence of oxidation kinetics on Cr content in bulk alloys can be explained by the values of k1 calculated based on the oxidation data reported by Rosecrans et al. for FeeCreNi alloys with various Cr contents in simulated PWR primary water environments [33,34]. The value of k1 increases and then decreases when Cr content increases from 5 to 30% (wt.), showing a k1 maximum at about 16% (wt.) Cr that is close to the Cr content of Alloy 600. Values of k1 for Alloy 600 were significantly lower than those for Alloy 690 in simulated PWR primary water at various dissolved hydrogen concentrations [23], based on the oxide film thickness data reported by P. Combrade et al. [23,35]. There are few data for the effects of Cr content in Ni-base alloys on the mechanical properties of oxide films formed in high temperature water. Rosecrans et al. [34] has reported that the oxide film rupture strain of NieFeeCr alloys increases with the Cr content. Increasing the Cr content from 5 to 30% (wt.), the film rupture strain in simulated PWR primary water with DH of 2e3 ppm at 288
C increases from about 8  104 to 2  103. Young et al. [8] have analyzed the chemical compositions at grain boundaries in an Alloy 600 base metal and its weld HAZ by transmission electron microscope (TEM). In the Alloy 600 base metal, Cr content at grain boundaries can be 5% lower than that in the matrix. In the Alloy 600 weld HAZ, Cr content at grain boundaries can be 9% lower than that in the matrix. Significant depletion of Cr at grain boundaries in the base metal and HAZ of an Alloy 600Alloy 182 weld has been demonstrated [10]. The effects of thermal treatment conditions at 700
C and 600
C on Cr-depletion at grain boundaries in Alloy 690 were investigated with TEM by Kai et al. [36]. The grain boundary Cr concentration in Alloy 690 after the treatments at 700
C for 1e200 h is higher than 20% (wt.). These results for grain boundary Cr distribution have been simulated by Yin et al. [37]. The grain boundary Cr concentration in Alloy 690TT HAZ is found to be still high as demonstrated by the results of SSRTAES measurements, as shown in Fig. 12. The metallurgical and mechanical characteristics in a weld HAZ, as shown in Fig. 1, would affect the SCC behavior in high temperP ature water. The lower fraction of 3 boundaries in the HAZ would

decrease the SCC resistance, as shown in previous results for Alloy 600 HAZ and cold worked Alloy 600 [9,10]. The enhancing effect of strain-hardening on PWSCC of Alloy 600 has been reported [9,10,38]. In some cases the properties of grain boundaries and cold work are correlated. Cold work can increase the hardness, the strength and the fraction of random boundaries. The significant effects of strain-hardening on PWSCC of Alloys 600 and 690 have been reported [1,3,9e16,38,39]. Andresen et al. [12,15] has reported that crack growth rates in most cases of CGR tests of cold worked Alloy 690 were in the range of 1012e1011 m/ s, but the CGRs of 1-dimensionally cold rolled Alloy 690 tested in the S-L orientation can be as high as 4  1010 m/s. The enhancing effect of cold work in SCC growth rates of Alloy 690 were reported in other works [11e16,39]. The local compositions at grain boundaries in Ni-base alloys could have a significant effect on intergranular SCC in high temperature water environments. The effect of Cr-depletion at grain boundaries for SCC of sensitized stainless steels in boiling water reactors was reported [20,21,40]. Experimental data show that Cr depletion at grain boundaries plays an important role in SCC of sensitized stainless steel [20,21,40] and Ni-base alloys in oxygenated high temperature water environments [41]. The difference of grain boundary Cr concentration could be one of the contributing factors for the difference in SCC behaviors of Alloy 600 HAZ and Alloy 690 HAZ. The beneficial effect of high bulk Cr concentration in Ni-base weld metals on PWSCC has been reported by Morton et al. [26]. These results are consistent with the modeling of intergranular SCC with Eq. (10) and available data related to the effect of Cr content on oxidation rate constant and film degradation strain. Bruemmer et al. [42] have pointed out the importance of grain boundary carbide coverage in Ni-base alloys on PWSCC susceptibility. Arioka et al. [39] have pointed out that PWSCC susceptibility of Alloy 690 decreases with increasing grain boundary Cr concentration but increases with increasing grain boundary carbide coverage. Based on the results of microstructure, local deformation P and microchemistry in Alloy 690 HAZ, lower fraction of 3 and high strength would enhance the PWSCC growth. Higher Cr concentrations in matrix and at grain boundaries in Alloy 690TT HAZ than those in Alloy 600 HAZ would contribute the lower PWSCC susceptibility in Alloy 690 HAZ than that in Alloy 600 HAZ, and the lower PWSCC growth rate in cold worked Alloy 690TT than those in Alloy 600.

Z. Lu et al. / Journal of Nuclear Materials 465 (2015) 471e481

4. Conclusions Distribution of HV, grain boundary chemistry, type of grain boundary and fraction, local deformation, and PWSCC behavior of P an Alloy 690TT HAZ are measured. Hardness and fractions of 3 boundaries increase and KAM decreases with increasing distance from the weld fusion line. The metallurgical and mechanical inhomogeneity in the Alloy 690 HAZ is similar to that in an Alloy 600 HAZ. Grain boundary Cr depletion in the Alloy 690 HAZ is less significant than that in the Alloy 600 HAZ. Much lower PWSCC growth rates are observed in the Alloy 690 HAZ and 40% CR Alloy 690 than those in the Alloy 600 HAZ and 40% CR Alloy 600, respectively. Cold wok significantly increases the SCC sensitivity of Alloy 690TT. Cr content in the alloys is thought to be one of the causes for the difference in PWSCC behavior of Alloy 690 HAZ and Alloy 600MA HAZ. High Cr contents in the bulk and at grain boundaries of Alloy 690 HAZ significantly decreases the SCC growth rate, which makes the effects of other factors such as hardening and P the fraction of 3 boundaries less significant. Acknowledgments

[14]

[15]

[16]

[17] [18]

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This work has been partly supported by PEAC-E and POLIM programs. Z.P. Lu acknowledges the support of the Ph.D. Programs Foundation of Ministry of Education of China No.20123108110021 and the International Cooperative Project sponsored by Science and Technology Commission of Shanghai Municipality No. 13520721200.

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