Stress corrosion cracking of simulated heat-affected zone in a CF8A weld in high temperature water

Journal of Nuclear Materials 527 (2019) 151810

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

Stress corrosion cracking of simulated heat-affected zone in a CF8A weld in high temperature water Tai-Cheng Chen a, Hau-Han Chang b, Jiunn-Yuan Huang a, Leu-Wen Tsay b, * a b

Nuclear Fuels and Materials Division, Institute of Nuclear Energy Research, Taoyuan, 32546, Taiwan, ROC Institute of Materials Engineering, National Taiwan Ocean University, Keelung, 20224, Taiwan, ROC

h i g h l i g h t s
Fast cooling after thermal simulation increased the ferrite content of CF8A.
Thermal simulation obviously improved the impact toughness of the high ferrite CF8A.
Cracks were found to be more likely to propagate along the d/g interfaces of CF8A.
Premature cracking of the g phase occurred in the low ferrite CF8A in high temperature water.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 19 April 2019 Received in revised form 18 September 2019 Accepted 18 September 2019 Available online 18 September 2019

In this study, thermal simulation was performed to conduct thermal cycles on CF8A cast austenitic stainless steel substrates with different ferrite contents to simulate the microstructures that form near a weld fusion boundary. The effects of ferrite content and morphology on the mechanical properties of CF8A were investigated. The stress corrosion cracking (SCC) susceptibility of the specimens was evaluated with slow strain rate tensile (SSRT) tests in a simulated boiling water reactor (BWR) coolant environment. The results indicated that an increase in the ferrite content of the CF8A base metal led to a minor increase in hardness. Thermal simulation caused a marked increase in ferrite content, especially in the high ferrite CF8A substrate. The yield strengths of the tested samples increased obviously with increasing ferrite content, but the ductility decreased. The ultimate tensile strengths of the samples reached a plateau when the ferrite number (FN) was greater than 18. The applied thermal simulation greatly improved the impact toughness of high ferrite CF8A in comparison with the original substrate. The results indicated that the samples with high ferrite content had a shorter rupture life but greater fracture strength than those with low ferrite content strained in high temperature water. Moreover, cleavage-like fracture was found in all of the samples that suffered from SCC. The SCC cracks were observed to be more likely to propagate along the d/g interfaces of a CF8A. Furthermore, crack initiation and growth inside the g of low ferrite CF8A contributed to a great loss in strength in a high temperature water environment. © 2019 Published by Elsevier B.V.

Keywords: CF8A stainless steel Thermal simulation Heat-affected zone Stress corrosion cracking Boiling water reactor

1. Introduction The thermo-mechanical property is an important index for understanding the deformation behavior of an alloy during hot/ cold working, welding or service. The changes in mechanical properties resulting from a specific thermal cycle can be evaluated by thermo-mechanical simulations, which incorporate the

* Corresponding author. E-mail address: [email protected] (L.-W. Tsay). https://doi.org/10.1016/j.jnucmat.2019.151810 0022-3115/© 2019 Published by Elsevier B.V.

application of heating/cooling cycles and tensile/compressive load simultaneously. The hot deformation behavior and processing maps of titanium alloys [1,2] and Al alloys [3] have been constructed by using thermal-mechanical simulation. Through the thermal simulation, different welding thermal cycles, from a single pass to triple passes, have been investigated for the pitting corrosion resistance of 2304 duplex stainless steel (SS) [4]. With the use of thermal simulation, liquation cracking in the simulated heataffected zone (HAZ) of Ni-based superalloy welds is found to be closely related to the precipitation of gamma prime [5], carbide [6], boride and intermetallics [7,8], which is deemed to account for the

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poor weldability of Ni-based superalloy. Cast austenitic stainless steels (CASSs) consisting of austenite (g) and ferrite (d) at room temperature are widely used in nuclear power plants for components such as pump and valve casings, fittings and primary coolant pipes [9,10]. Dissimilar metal welding is widely used for joining low alloy steels to SS components in the nuclear power industry [11]. Alloys 82 and 182 are often used as filler metals to join the safe end (austenitic SS) and the A508 (low alloy steel) feed-water nozzle [11]. Moreover, the occurrence of primary water stress corrosion cracking (PWSCC) has been observed in several dissimilar welds of pressurized water reactors (PWR) [12]. Preemptive weld overlays have been applied to the piping system in a nuclear power plant to mitigate the occurrence of stress corrosion cracking (SCC). The introduction of residual compressive stress to the inner surface tends to reduce both the SCC susceptibility [13] and the potential for fatigue crack growth of the pipes [14]. In practical application, to avoid hot cracking of the Ni-based overlay weld, a 308 SS buffer layer is applied before an IN52 M overlay is deposited onto CASS [15e18]. Moreover, welding austenitic SSs can result in an obvious increase in ferrite content in the weld metal (WM) and in the HAZ near the fusion boundary (FB), as compared with that in the base metal (BM) [19]. In the case of a dissimilar weld, the great changes in microstructures in the HAZ near the FB of a CASS weld will cause obvious variation in the mechanical properties and/or the SCC resistance, which are matters of great concern in a nuclear power plant. In the open literatures, CASSs have been reported to be susceptible to SCC in simulated reactor coolant environments [20e22]. However, SCC in the HAZ near the FB of a CF8A weld had not been studied before. Therefore, the objective of this study is to investigate the effect of ferrite content and morphology in the simulated HAZ on the SCC susceptibility of a CF8A weld in simulated reactor coolant environment. With a Gleeble simulator, the effects of welding thermal cycles on the microstructures, tensile properties and impact toughness of CF8A with different ferrite contents were investigated in this study. Slow strain rate tensile (SSRT) tests were performed to evaluate the SCC susceptibility of various specimens in a simulated boiling water reactor (BWR) coolant environment. The fracture morphology was examined with a scanning electron microscope (SEM), and the phase changes adjacent to the fracture surface were identified by electron backscatter diffraction (EBSD). It is expected that a great change in ferrite content and morphology in the HAZ adjacent to the FB of a CASS weld will play an important role in SCC susceptibility.

either 8.0 or 21.4 were used in this work, and they were signified as the low ferrite (LF) and high ferrite (HF), respectively. During welding, the microstructures and mechanical properties of the narrow HAZ are obviously different from those of other regions in a weld. Similar microstructures formed in the HAZ of a weld can be accessed by thermal simulation. The temperature of the HAZ near the FB in a CASS is expected to peak slightly below its melting temperature, followed by cooling down by the chill of the substrate. Fig. 1 shows the dimensions of the samples for thermal simulation (Fig. 1a) and tensile tests (Fig. 1b). The thickness of the tested specimens was 2.5 mm. A Gleeble 1500 simulator was used to generate microstructures similar to those formed in the HAZ near the FB of a CF8A weld. The samples were heated by the Gleeble simulator from room temperature to 1180  C at a heating rate of 30  C/sec, followed by heating to the peak temperature of 1320  C at a heating rate of 15  C/sec. To achieve ample g to d transformation, the tested samples were held at 1320  C for 1 min before being chilled by water spray. During thermal simulation, the tested samples were not constrained by external fixtures. The LF and HF samples after thermal simulation were denominated as LG and HG, respectively. The FN of the LG sample was 18.3, and that of the HG sample was 46.9. 2.2. Mechanical and SCC tests The hardness of different samples was measured using a Mitutoyo MVK-G1500 microhardness tester with an applied load of 300 gf and a dwell time of 15 s. A Charpy V-notch (CVN) impact specimen with a thickness of 2.5 mm, which met the ASTM E23 specification, was used to determine the impact toughnesses of various specimens at room temperature. The specimen dimensions for the tensile and SSRT tests are shown in Fig. 1b. To evaluate the effect of thermal simulation on the mechanical properties of CF8A, the tested samples were wire-cut from the simulated samples, as shown in Fig. 1a. Tensile tests at a strain rate of 6  104/sec in air at room temperature and SSRT tests at a strain rate of 1  106/sec in high temperature water were performed. A strain gage of 10 mm gage length was adhered to the tensile sample to measure the elongation in air. For the SSRT tests, the specimens were installed in the Toshin Kogyo SCC testing system, which is a closed-loop, servoelectric machine with an integrated water circulation loop. The details of the simulated BWR water chemistry for the SSRT tests are listed in Table 2. All the mechanical properties shown in this work are the average of at least four tests for each sample.

2. Material and experimental procedures

2.3. Microstructural examination

2.1. Material

Samples for microstructural examination were obtained through proper metallographic sample preparation process. The testing samples were carefully sliced by a low-speed cutter and hotmounted in the resin, after that, the mounted samples were grounded with grinding papers up to 2000 grit and then polished by using Al-oxide powder. The microstructures of various

The chemical compositions of the centrifugally cast CF8A used in this study are listed in Table 1. The ferrite contents of various samples were determined with a Fischer FMP30 ferritescope and evaluated by the ferrite number (FN). Two heats of CF8A with FN of

Table 1 Chemical compositions of CF8A used in this study. Specimen

LF (FN: 8.0) HF (FN: 21.4)

Chemical composition (wt.%) C

Cr

Ni

Si

Mn

P

S

Fe

0.054 0.066

19.42 21.95

8.50 8.28

0.34 1.15

0.21 0.88

0.039 0.028

0.004 0.008

bal. bal.

Note: [23,24]. Creq ¼ Cr þ Mo þ 1.5 Si þ 0.5 Nb. Nieq ¼ Ni þ 30 C þ 0.5 Mn.

Creq

Nieq

Creq/Nieq

19.9 23.7

10.2 10.7

2.0 2.2

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Fig. 1. Schematics of the samples for (a) thermal simulation and (b) tensile tests.

3. Results

Table 2 Water chemistry of the SSRT tests. Test parameters

Value

Pressure (MPa) Temperature ( C) Inlet/outlet conductivity (mS/cm) Inlet/outlet oxygen (ppm) Inlet/outlet hydrogen (ppb) Inlet/outlet pH

10 300 0.058/0.091 7.85/7.13 0.1/0.2 6.96/6.15

specimens were inspected with an Olympus BX51 optical microscope (OM) and a JEOL JSM-7100F field emission SEM. The yielding volume of an electron beam will affect the compositional measurements at a specific site. To avoid the interference of background, a JEOL JSM-7100F SEM equipped with an Oxford Instruments X-MaxN energy-dispersive spectrometer (EDS) was used first to check the compositions of the d and g phases at different sites. The chemical compositions of various phases at specific sites, which had been checked previously by EDS, were further determined by JEOL JXA-8200 superprobe electron probe micro-analyzer (EPMA). The spot size is about 1 mm and operating at 15 kV. The shown compositions of the d or g phases were the average of at least three measurements. Moreover, the specimens were also examined with a JEOL JSM-7100F SEM equipped with an Oxford Instruments NordlysMax2 EBSD detector to identify distinct phases in the samples. Moreover, the strain distribution in the tested samples can be analyzed by inputting the EBSD map into the HKL-Channel 5 software for post-processing of data. The fracture features of the tested samples were examined with a field emission SEM.

3.1. Microstructural observation Type 308L SS rods of 1.2 mm diameter were used as the filler to be deposited on the two CF8A substrates (LF and HF) by gas tungsten arc welding in multiple passes. The welding parameters were as follows: welding current of 185 A, welding voltage of 19 V, fillerfeeding speed of 1050 mm/min, and welding speed of 125 mm/min. The SEM micrographs and EBSD phase maps around the FB of the overlay welds, which included the WM and HAZ, are shown in Fig. 2. The results indicated that the FN of the WM determined by ferritescope was about 10. As shown in Fig. 2a and b, the d ferrite that formed in the WM was much finer in size than that in the CF8A substrates. It was deduced that much faster cooling after welding resulted in the formation of fine d ferrite in the WM, as compared with that in the centrifugal cast substrates. Moreover, the d ferrite in the WM exhibited epitaxial growth from the ferrite in the CF8A substrates (Fig. 2c and d). The d ferrite formed in the HAZ near the FB of the HF weld (Fig. 2c) was much greater in amount and coarser in size than that of the LF weld (Fig. 2d). Generally, the g phase was the stable and major phase at room temperature in the CASSs. Near the FB, the d content of the LF sample (Fig. 2d) was lower than that of the HF sample (Fig. 2c). According to image analysis, the estimated d content in the HAZ near the FB of the HF sample was about 42% (g: 58%), which was obviously higher than that of its substrate. This result confirmed that the welding thermal cycle would greatly increase the ferrite content in the HAZ near the FB of a CASS weld. Fig. 3 presents optical micrographs showing the microstructures of various specimens. As shown in Fig. 3a, the LF specimen with FN of 8.0 consisted of thin, island-chain-like d ferrites dispersed in the

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Fig. 2. SEM micrographs of the (a) HF, (b) LF samples and EBSD phase maps of the (c) HF and (d) LF samples around the fusion boundary of the overlay welds.

Fig. 3. EBSD phase maps of the (a) LF, (b) HF, (c) LG and (d) HG specimens.

g matrix. As listed in Table 1, increasing the Creq/Nieq of the CF8A from 2.0 to 2.2 would cause an obvious increase in the FN from 8.0 to 21.4. The results indicated that a small increase in the Creq/Nieq ratio would cause a great increase in the FN of CF8A. The d ferrite morphology of the HF sample (Fig. 3b) was similar to that of the LF

sample; the difference between them was that the d ferrites were coarser and their amounts were greater in the latter. When Creq/ Nieq is greater than 1.95 [23,24], d ferrite will be the primary phase after solidification at elevated temperature. Moreover, a rapid cooling rate limits the time available for the d to g phase

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transformation, and helps retain more ferrite in the CASS. The d ferrite morphology and distribution in the LG specimen are displayed in Fig. 3c. A moderate increase in FN from 8.0 to 18.3 was found in the LF substrate subjected to thermal simulation. The results indicated that thermal simulation also led to a great rise in the FN of the high ferrite CF8A. The FNs of the HF and HG samples were 21.4 and 46.9, respectively. It was obvious that imposing thermal simulation greatly increased the ferrite content of CF8A, especially that of the high ferrite sample. It was seen that the island-like d ferrites in the LG sample were wider and larger than those in the LF sample. The results also indicated that the FN of the LG sample (18.3) was only slightly lower than that of the HF (21.4). As compared with that in the HF sample, the d ferrite morphology in the LG sample was more likely to be isolated and coarser. The FN of the HG sample was as high as 46.9, and irregular d ferrites formed extensively in the sample (Fig. 3d).

3.2. EPMA analysis As listed in Table 1, Cr and Ni are the major elements of CF8A, and the compositions can also be related to the solidified microstructures. Fig. 4 presents the changes in the Cr and Ni contents of

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the g (Fig. 4a) and d (Fig. 4b) phases in distinct CF8A substrates before and after thermal simulation, as measured by EPMA. As listed in Table 1, the HF sample had higher Cr but lower Ni concentrations than the LF sample did. According to the EPMA results, the g was rich in Ni but lean in Cr relative to the d ferrite. Moreover, the range of Ni content in the g phase (8.3e9.8 wt%, Fig. 4a) in the four different samples was relatively limited as compared to that in the d phase (3.6e6.3 wt%, Fig. 4b). In each group, thermal simulation increased the ferrite content, which led to an increase in Ni content in the g phase in response. The LG sample had the highest Ni content (9.8 wt%) in the g of all the samples, as shown in Fig. 4a. The Ni content in the d ferrite increased in the order of LF < HF < LG < HG samples. It was noticed that the Ni content (6.3 wt%) in the d ferrite of the HG sample was higher than that in other samples, as shown in Fig. 4b. It was expected that the g in the CASS would transform into d and cause a rise in the Ni content of the d ferrite during thermal simulation. Upon rapid chilling to room temperature, the g grew and replaced the d as the major and stable phase. The high Ni content in the retained d ferrite of the simulated sample could be due to the limited Ni diffusion during quenching. Therefore, the Ni content of the d ferrite of the simulated sample was higher than that of the counterpart sample. Thermal simulation increased the ferrite content in the simulated sample, which was accompanied by a decrease in Cr content (Fig. 4b). By contrast, the Cr content in the g phase increased slightly in the simulated sample, in comparison with the original substrate (Fig. 4a). The decline in g content in the simulated sample led to a rise in Cr concentration in the g phase. Similar results have been pointed out in investigating the effect of solution temperature on the changed chemical compositions of the d and g phases in a casted duplex stainless steel [25]. As a whole, a great change in the d to g ratio was anticipated to result in obvious variation in the mechanical properties. 3.3. Mechanical properties in air The microhardnesses of the LF, HF, LG and HG samples were HV 158, 164, 166 and 180 after repeated tests, respectively. Despite the great differences in FNs among the specimens, the results indicated that an increase in ferrite content could cause only a minor increase in hardness. Table 3 lists the mechanical properties of various samples tested in air at room temperature. The tensile tests showed that the yield strength of the test samples increased obviously with increasing ferrite content; however, the ductility varied in the reverse manner. The yield strength of the HG sample was twice as much as that of the LF sample. The ultimate tensile strengths of the HF, LG and HG samples were similar, about 600 MPa. It seemed that the ultimate tensile strength of the samples reached a plateau when the FN was greater than 18. As listed in Table 3, the sub-size Charpy impact energies of the LF, HF, LG, and HG samples were 40, 16, 40 and 32 J, respectively. These results indicated that the HF specimen had the lowest resistance to impact fracture. The LF and LG samples exhibited similar impact toughness, indicating that the thermal cycle imposed by the thermal simulation had little effect on the impact toughness of low ferrite CF8A. Furthermore, it was noticed that thermal simulation obviously improved the impact toughness of the high ferrite CF8A. As shown in Fig. 3, thermal simulation increased the ferrite content of the high ferrite CF8A. The discrepancy in impact toughness between the HF and HG samples could be attributed partly to the irregularly blocky ferrite inducing a tortuous fracture path. 3.4. SCC in high temperature water

Fig. 4. The results of EPMA analysis showing the changes in Cr and Ni contents in (a) g and (b) d phases for distinct CF8A substrates before and after thermal simulation.

Table 4 shows the typical results of the fracture strength,

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Table 3 The results of tensile and impact tests of various samples tested in air at room temperature. Specimen

Yield strength (MPa)

Ultimate tensile strength (MPa)

Elongation (%)

CVN impact toughness (J)

LF HF LG HG

195 ± 10 217 ± 16 274 ± 15 467 ± 21

584 ± 19 598 ± 7 598 ± 13 602 ± 14

53 ± 5 40 ± 3 46 ± 2 30 ± 4

40 ± 3 16 ± 2 40 ± 1 32 ± 1

Table 4 The properties of various samples subjected to SSRT tests in high temperature water. Specimen

Fracture strength (MPa)

Fracture life (h)

Elongation (%)

LF HF LG HG

355 ± 3 504 ± 1 389 ± 4 560 ± 12

156 ± 2 129 ± 2 133 ± 3 123 ± 1

11 ± 1 9±0 10 ± 1 9±1

fracture life and elongation of various specimens tested in high temperature water. The fracture lives of the samples increased in the following order: HG < HF < LG < LF (123, 129, 133, and 156 h, respectively). In addition, the fracture strengths of various samples in high temperature water decreased in the following order: HG > HF > LG > LF (560, 504, 389, and 355 MPa, respectively). These results indicated that the fracture lives and fracture strengths of the samples had opposite tendencies. After the SSRT tests in high temperature water, the strained samples with high FNs exhibited high fracture strengths before rupture. As compared with the ultimate tensile strength tested in air (Table 3), the loss in strength in high temperature water decreased in the following order: HG < HF < LG < LF (42, 94, 209, and 229 MPa, respectively). It seemed that in high temperature water, the LF and LG group suffered higher losses in strength than the HF and HG group did. The results also indicated that the tensile elongations of the HF, LG and HG samples after SSRT tests in high temperature water were similar, while the elongation of the LF sample was a little higher than those of the other samples. 3.5. SEM fractographs Fig. 5 displays the fracture features of distinct samples after tensile and impact tests in air. Ductile dimple fracture was found in the LF and LG samples after tensile-straining in air (Fig. 5a). Terracelike fracture with secondary cracks at the boundaries was found in the HF sample after tensile testing in air (Fig. 5b). Such brittle feature accounted for the low ductility of the HF sample relative to those of the LF and LG samples. The fracture appearance of the HG sample showed small facets with fine cracks along the interfaces (Fig. 5c). As compared with the HF sample, the HG sample had an increased extent of brittle feature, which was associated with the inferior tensile ductility. As listed in Table 3, the HF sample had the lowest Charpy impact energy of all the samples being tested. The impact fracture appearance of the HF sample displayed quasicleavage fracture (Fig. 5d). By contrast, ductile dimple fracture was found in the LF and LG samples after impact tests, whereas shallow dimples mixed with small facet fractures were obtained in the HG sample (not shown). Less cleavage fracture in HG sample meant cracks were less likely to propagate straight forward or the HG sample did not fracture as brittle as the HF sample. The difference in the impact fracture appearance between the HF and HG samples was consistent with their discrepancy in impact toughness. The typical fracture appearances of the samples tested in high temperature water are shown in Fig. 6. As can be seen in that figure, cleavage and/or quasi-cleavage with secondary cracks along the

interfaces were/was more likely to be observed in the embrittled areas of the LF and LG samples (Fig. 6a and b), especially in the LF sample. By contrast, small brittle facets and quasi-cleavage were found in the embrittled areas of the HF and HG samples (Fig. 6c and d). However, g-d interfacial separations were more often found in the embrittled zones of the HG sample (Fig. 6d-left). Numerous side cracks were found on the HG sample surface after SSRT tests (Fig. 6d-right), also implying the high cracking susceptibility of the sample. 3.6. EBSD identifications Fig. 7 presents the EBSD phase maps showing the phases around the fracture zones of the samples after tensile straining in air or high temperature water. It is known that metastable g will transform into martensite during straining [26e31], leading to an increase in hydrogen embrittlement susceptibility. In the samples strained in air, intense plastic deformation at the fracture zone caused the g to transform to a0 martensite, which had a crystal structure similar to that of the d ferrite (Fig. 7a). Slip bands, d ferrite and deformation-induced a0 martensite were observed at the fracture zone of the LF sample strained in air, indicating that extensive plastic deformation of the LF sample occurred before final fracture during straining in air. Moreover, cracks tended to initiate and propagate along the d/g interfaces of the HF sample (Fig. 7b). Deformation-induced a0 martensite was also less likely to be found at the fracture zone of the HF sample strained in air. The LG and HG samples strained in air behaved similarly to the LF and HF samples, respectively. In the LF samples strained in high temperature water, far fewer slip bands and no induced a0 martensite were found (Fig. 7c). The absence of induced a0 martensite adjacent to the crack tip meant that in high temperature water, cracking instead of plastic deformation occurred in the LF sample. The results also revealed that the SCC cracks of the LF samples were more likely to propagate along the d/g interfaces (Fig. 7c). Moreover, it was noticed that fine cracks could also initiate individually in the g matrix during straining in high temperature water, as shown in Fig. 7c. It was deduced that crack initiation and growth in the g might have contributed greatly to the high loss of strength for the low ferrite CF8A strained in high temperature water. As mentioned previously, in high temperature water, the HG sample with high ferrite content possessed a higher fracture strength and lower loss in strength, but shorter fracture time than those of other samples. The EBSD phase map (Fig. 7d) displays the initiation and growth of SCC cracks preferentially at the d/g interfaces. The increased ferrite contents of the CASS were associated with an increase in d/g interfaces. In high temperature water, the easy linkage of numerous interfacial cracks in the HG sample (Fig. 6d) resulted in a short fracture life. It was obvious that the d/g interfaces were also the weak path for the SCC cracks of all the CASSs in high temperature water. The strain distribution in the tested samples can be analyzed by inputting the EBSD map into the HKL-Channel 5 software for data post-processing. The corresponding strain contours of those samples investigated in Fig. 7 are shown in Fig. 8. The locations with intense slip bands (Fig. 7a), possibly consisting of embedded a0

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Fig. 5. Tensile fracture features of the (a) LF, (b) HF and (c) HG samples tested in air; (d) fracture features of the HF sample after Charpy impact test.

Fig. 6. Fracture features of the (a) LF, (b) LG, (c) HF and (d) HG samples after SSRT tests in high temperature water. Note that images (a) to (c) and (d)- right were taken from the fracture surface; while image (d)- left was taken from the side of the failed sample.

martensite in the red-colored zones, were consistent with relatively high strain in the LF sample strained in air (Fig. 8a). As shown in Fig. 7b, cracks were found to initiate and propagate along the d/g interfaces of the HF sample strained in air. As a whole, the strain of the elongated d ferrite was at the same level as that of the adjacent g matrix (Fig. 8b); both phases were colored green. This revealed

the fact that the loading strain was spread around the fracture zone in the HF sample strained in air. After being strained in high temperature water, the g matrix (Fig. 8c and d) exhibited low strain (in blue). The low g strain around fracture zone implied plastic deformation was less likely to occur while straining in high temperature water. As shown in Fig. 8c, those fine g debris that adhered

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Fig. 7. EBSD phase maps around the fracture zone of the (a) LF and (b) HF samples after tensile tests in air; (c) LF and (d) HF samples after SSRT tests in high temperature water.

Fig. 8. EBSD strain distribution maps around the fracture zone of the (a) LF and (b) HF samples after tensile tests in air; (c) LF and (d) HF samples after SSRT tests in high temperature water.

to the fracture surface of the sample had low fracture strain (in blue), which was related to its brittle nature in high temperature water. Moreover, extensive spreading of the strain did not occur in the fracture zones of the samples strained in high temperature water. In fact, strain localization was observed in some of the ferrite

(in green) in Fig. 8c and d. Concentrated strain around the ferrite would facilitate the occurrence of cracking at the d/g interfaces. Therefore, the initiation and growth of microcracks in the g phase and/or at the d/g interfaces was responsible for the SCC cracking of the CASS in high temperature water.

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4. Discussion Generally, the ferrite possesses poorer toughness and ductility than the austenitethus, an increase in the ferrite contents of the CASSs reduces the ductility and toughness. After the thermal simulation, the FN of the HG samples increased greatly to 46.9, as opposed to 21.7 in the HF sample. The results of impact tests indicated that the impact toughness of the HG samples was greater than that of the HF samples. As shown inFig. 3, skeletal ferrite formed in the HF sample (Fig. 3b), whereas blocky and irregular island-like ferrites were found in the HG sample (Fig. 3d). After the thermal simulation of high ferrite CF8A, the improved d ferrite morphology had a lowered stress concentration as compared with that of the as-cast substrate; thus, the impact toughness was improved. It was also reported that impact energies of a casted duplex stainless steel increased with increasing solution temperature from 980 to 1380  C [25]. The sample with the same ferrite and austenite fraction had the best impact properties [25]. By contrast, the same improvement in impact toughness by thermal simulation was not obtained in the low ferrite CF8A. The exact reasons for such results were not realized at this moment. The ferrite occupied a minor portion (FN ¼ 8) in the LF sample, thus, the LF sample showed high impact toughness (40 J). It was deduced that the harmful effect of high ferrite content in the LG sample relative to the LF one, was mitigated by its beneficial ferrite profile of the former. Thus, the LG and LF samples had the same resistance to impact fracture. To evaluate the stability of austenitic SSs, the temperature at which 50% of the austenite phase transforms into martensite during tensile testing at a true strain of 0.3 is defined as the Md30 temperature [32]. The formula that relates the Md30 to the chemical composition of an alloy is as follows: Md30 (ºC) ¼ 413e462 (%Cþ%N) - 9.2 (%Si) - 8.1 (%Mn)-13.7(%Cr) - 9.5(%Ni) - 18.5 (%Mo). The lower the Md30 temperature, the more stable the g will be. Fig. 4 displays the Cr and Ni contents in the d and g phases in different samples. Regarding the HF and HG samples, thermal simulation lowered the differences in alloy concentrations of the g and d phases. In other words, the chemical compositions of the different phases became more uniform. In the HF sample, the Ni contents of g and d were 8.3 and 4.7 wt%, respectively; by contrast, they were 9.1 and 6.3 wt% in the HG sample. The Cr contents of the g and d phases in the HF sample were respectively 20.6 and 26.7 wt%, in contrast to 22.0 and 25.2 wt% in the HG sample. After thermal simulation, an increase in the Ni contents of the g and d phases might partially explain the improved toughness of both. According to the Md30 temperature equation, an increase in the Cr content of the g phase was also helpful to increase the g stability. Therefore, the combination of altered d ferrite morphologies and lowered differences in composition of the g and d phases led to significant improvement in the impact toughness of the HG sample with respect to that of the HF sample. The combination of ferrite and austenite in duplex SSs, i.e., 2205 and 2507, has been confirmed to provide SCC resistance superior to those of 304 and 316 SSs in chloride-containing environments [33e36]. In this work, an increase in FN was associated with an obvious increase in yield strength but accompanied with a decrease in the tensile ductility of the CF8A in air at room temperature. Cleavage-like fracture (Fig. 5b) and interfacial separation along the g/d boundary (Fig. 7b) confirmed the low ductility of the HF sample relative to that of the LF sample. The low ductility of the HF samples was associated with the slip characteristics and the g/d ratios in the CF8A substrates. The lower g content in the high ferrite CF8A would reduce the occurrence of possible slips or the motion of dislocations in the g, in comparison with those in the low ferrite CF8A. Moreover, the extensive slips in the g would be restricted with an

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increase in the d content in the CF8A. Therefore, an increase in FN was responsible for the increased yield strength of the CF8A in air at room temperature. The water chemistry, listed in Table 2, seemed to have a very mild effect on the CASSs. Previous studies have shown that CASSs are sensitive to SCC in simulated PWR primary water [20,22,37,38]. The brittle fracture of ferrite on the surface of a CASS is a prerequisite for SCC in simulated PWR water [20]. In this work, the strained samples with high FNs exhibited high fracture strength and low loss in strength, but short time-to-fracture, in high temperature water under SSRT tests, especially the HG sample. In general, the loss in strength is often used as an index of the SCC susceptibility of a tested sample. However, the low loss in strength of the HG sample in high temperature water might not properly be related to low SCC susceptibility. The ease of linkage of interfacial cracks in the HG sample could be one of the causes of the short time-to-fracture in high temperature water relative to those of other samples under SSRT tests. It has been reported that a high ferrite CASS is more sensitive to ductility loss than a low ferrite one, as determined by SSRT tests in high temperature water [22]. As shown in Fig. 6d, the side surface of the HG sample was covered with numerous fine cracks, which were responsible for the low ductility and short fracture time in high temperature water. Many interfacial separations also provided evidence of high cracking susceptibility of the sample. The d/g interfaces were found to be the weak path for the SCC cracks in all the samples, indicating that d/g interface likely played an important role in cracking in high temperature water. As shown in Fig. 7c, fine cracks initiated in the g matrix of the LF sample strained in high temperature water. The absence of deformation induced a0 martensite around the crack tip also meant that cracking instead of plastic deformation of the g phase was more likely to occur in the LF sample affected by SCC in high temperature water. Although the SCC cracks were found to be more likely to propagate along the d/g interfaces of all samples (Fig. 7), crack initiation and growth in the g phase of the low ferrite CF8A might also have contributed to the high loss in strength in high temperature water. Those fine g debris adhered to the fracture surface was of low fracture strain (Fig. 8c), showing their brittle nature in high temperature water. As compared to the g phase strained in air, the low g strain around fracture zone was consistent with the absence of dense slip lines therein, while straining in high temperature water, as shown in Fig. 8c and d. The exact reasons for the changed role of slips and the embrittlement of the g phase in high temperature water will require further studies. In this work, the testing cycle for each sample in high temperature water was about one week long, which was much shorter than the actual service life of a dissimilar weld in a nuclear power plant. Under the SSRT tests, all the tested samples were subjected to high loading before final cracking. With low ferrite content, the microcracks tended to initiate in the coarse g phase, resulting in showing high SCC susceptibility for the low ferrite CF8A tested in high temperature water. Furthermore, the dislocation motions in the g could be retarded by the d, thus, pile-up of dislocations was more likely to occur at the d/g interfaces. As mentioned previously, the SCC cracks had a trend to initiate and propagate along the d/g interfaces of all samples. Therefore, increasing the d/g interfaces in the CF8A would make the linking of SCC cracks more easily, showing lots of interfacial separations. Such results were seen for the HG samples tested in high temperature water. However, the SCC susceptibility of the CASS might change if low loading was applied. While corrosion or pitting corrosion at the d/g interfaces were the major causes for the crack initiation, the samples with high ferrite content would become susceptible to SCC in a simulated boiling water reactor (BWR) coolant environment.

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5. Conclusions 1. An increase in ferrite content in distinct CF8A samples resulted in minor increases in hardness. Thermal simulation of heating the samples to 1320  C and cooling them by water spray caused marked increases in ferrite content, especially in the high ferrite CF8A substrate. The FN of the high ferrite CF8A substrate after thermal simulation was about 47. The ferrite morphology changed from a skeletal structure in low ferrite CF8A to blocky and irregular island-like ferrite in high ferrite CF8A. 2. The yield strengths of the test samples increased obviously with increasing ferrite content; however, the tensile elongation varied in the reverse manner. The presence of d ferrite effectively impeded the occurrence of extensive slips in the g of the CF8A during straining. The Charpy impact energies of the LF, HF, LG, and HG specimens were 40, 16, 40 and 32 J, respectively. Thermal simulation obviously improved the impact toughness of the high ferrite CF8A. 3. The HF and HG samples exhibited high fracture strength and low loss in strength, but short time-to-fracture, in high temperature water during SSRT tests, especially the HG sample. The inherent low ductility of the HG sample might have been one of the causes of the short fracture time in high temperature water relative to those of other samples. Although the cracks were found to be more likely to propagate along the d/g interfaces, crack initiation and growth in the g phase of the low ferrite CF8A sample might also have contributed to the high loss in strength in high temperature water. 4. Terrace-like fracture with secondary cracks at the boundaries of the HF sample accounted for its low tensile ductility relative to those of the LF and LG samples in air. The tensile fracture appearance of the HG sample in air showed small facets with fine cracks along the d/g interfaces. Cleavage and/or quasicleavage with secondary cracks at the d/g boundaries were/ was more likely to be found in the samples strained in high temperature water than in those strained in air. The EBSD phase maps showed the absence of intense slip bands and induced a0 martensite around the crack tip, implying premature cracking instead of plastic deformation of the g phase in the CF8A affected by SCC in high temperature water.

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