Microstructure evolution and pitting corrosion behavior of UNS S32750 super duplex stainless steel welds after short-time heat treatment

Corrosion Science 121 (2017) 22–31

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Microstructure evolution and pitting corrosion behavior of UNS S32750 super duplex stainless steel welds after short-time heat treatment Ziying Zhang a,∗ , Hui Zhao a , Huizhen Zhang b , Jun Hu a , Jiarui Jin a a b

School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, China School of Management, University of Shanghai for Science and Technology, Shanghai 200093, China

a r t i c l e

i n f o

Article history: Received 2 August 2016 Received in revised form 3 February 2017 Accepted 4 February 2017 Available online 8 February 2017 Keywords: Super duplex stainless steel Welding Heat treatment Pitting corrosion

a b s t r a c t The effects of short-time heat treatment at different temperatures on the microstructure evolution and pitting corrosion behavior of UNS S32750 super duplex stainless steel welds were investigated. The results demonstrated that the ferrite phase in the as-welded heat-affected zone overwhelmed the austenite phase in content and was easily attacked by pitting corrosion. After short-time heat treatment, the excessive ferritization in the heat-affected zone was significantly alleviated. The highest pitting corrosion resistance of the welds was obtained at 1080 ◦ C. © 2017 Published by Elsevier Ltd.

1. Introduction Duplex stainless steels (DSSs) are widely used as structural materials in the chemical, petrochemical, marine, nuclear, and paper industries because of the attractive combination of their mechanical properties, weldability, and corrosion resistance in various types of environments [1–6]. Generally, the maintenance of roughly equal amounts of the ferrite and austenite phases, as well as the absence of the secondary phases, is mandatory to acquire good corrosion resistance and high strength in the alloys [7,8]. However, welding tends to disturb this phase balance [9–11]. Rapid heating and cooling cycles in DSS welds generally result in the excessive ferritization of the heat-affected zone (HAZ) and weld metal (WM) [10–12]. Furthermore, certain unwanted phases such as carbide, the sigma phase, the chi phase, and chromium nitrides are prone to form during welding [13,14]. The undesirable excessive ferritization and unwanted precipitations worsen the corrosion resistance of DSS welds [15,16]. The effects of the metallurgical factors, welding parameters, and welding processes on the microstructure evolution of DSSs welds were thoroughly investigated to enhance the corrosion resistance of the DSSs components [17–21]. Muthupandi et al. [17] reported that the excessive ferriti-

∗ Corresponding author. E-mail address: [email protected] (Z. Zhang). http://dx.doi.org/10.1016/j.corsci.2017.02.006 0010-938X/© 2017 Published by Elsevier Ltd.

zation of DSS welds is predicted to be alleviated by an enhancement of nickel content in the weld filler material. Jiang et al. [18] investigated the influence of Creq /Nieq on the pitting corrosion resistance and mechanical properties of UNS S32304 DSS welded joints with the thermal-mechanical simulator (Gleeble 3800) and discovered the austenite content in the high temperature HAZ increased with a decrease in the Creq /Nieq value. Kim et al. [19] reported that the austenite content in the WM and HAZ is enhanced effectively by combining Ar and N2 as the shielding gas during welding. Furthermore, the ferrite/austenite ratio also depends on the energy input for welding [9,20,21]. Although an optimal phase balance may be achieved in the WM of DSSs welds by adjusting the heat input to 0.3–1.5 kJ/mm [22], the austenite phase in the HAZ is still overwhelmed by the ferrite phase [23]. Several studies have focused on the post-weld heat treatment (PWHT) of DSSs welds to obtain an optimal phase balance [24,25]. Although PWHT effectively eliminates the excessive ferritization in DSS welds, it is impractical for larger weldments. Therefore, an on-line solution treatment was recommended in the workshop to meet the requirements of the practical application. The heat treatment technique generally minimizes the heat treatment time. Super duplex stainless steels (SDSS) were recently developed for the marine and petrochemical industries due to its excellent corrosion resistance in chloride and sulfide environments [26–28]. The pitting resistance equivalent (PRE) number of these alloys are reportedly greater than 40% [28]. However, the excessive ferriti-

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Table 1 Chemical compositions of UNS S32750 super duplex stainless steel. Element

C

Si

Mn

P

S

Cr

Ni

Mo

N

Fe

Wt%

0.016

0.40

0.75

0.021

0.001

24.88

6.88

3.79

0.269

Bal

Fig. 1. Schematic diagram of electrochemical sample. Table 2 Tungsten inert gas welding parameters. Welding current (A) Welding voltage (V) Welding speed (cm/min) Shielding gas flow: Ar (L/min) Backing gas flow: Ar (L/min)

160 15 2.7 18 5

zation of SDSS welds is heavier given their alloying with higher levels of chromium and molybdenum. Furthermore, tungsten inert gas (TIG) welding is prevalent in industry due to its comparatively easier applicability and better economic solutions. Therefore, investigating the effect of a short-time heat treatment on the microstructure evolution and pitting corrosion behavior of UNS S32750 SDSS welds can guide the reduction of UNS S32750 SDSS component security risks and initiate the innovative design of an alloy that exhibits excellent weldability. However, at this time, reports have not been published with regards to this issue. The present study characterized the effects of short-time heat treatment on the microstructure evolution and pitting corrosion behavior of UNS S32750 SDSS welds by morphological observation and electrochemical detection. The TIG welding process was executed with appropriate welding parameters to which the welds were expected to exhibit less excessive ferritization and no secondary phases. The relationships among the microstructure evolution, heat treatment temperature, and pitting corrosion behavior of the UNS S32750 welds were discussed in detail. An optimum short-time heat treatment temperature was determined. 2. Experimental

Fig. 2. Optical micrographs of different zones of the as-welded after etching in 30 wt% KOH electrolyte: (a) BM, (b) HAZ (between the lines), (c) WM.

2.1. Materials and heat treatment The base metals employed in this work is a commercial UNS S32750 SDSS produced by Outokumpu Stainless. The alloy has been cold rolled into plates of about 5 mm thickness and then annealed at 1100 ◦ C for 30 min in argon flow. The chemical compositions of base metals are shown in Table 1. The plates were butt welded with single TIG welding. The welding parameters are listed in Table 2. The heat treatment was performed in a muffle furnace at 1020, 1050, 1080 and 1100 ◦ C for a holding time of 3 min, followed by water quenching. These heat treatments were performed in the

“precipitation-free” temperature range for duplex stainless steels. In order to have the WM, HAZ and base metal (BM) included in the coupons, the specimens were cut into plates with a dimension of 20 mm × 20 mm. The specimens were successively ground to 1200 grit using SiC abrasive papers, and then polished with diamond paste to 0.25 ␮m.

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Table 3 Mean chemical compositions of the test specimens (wt%) after twelve times EDS measurements. Sample

Zone

Ferrite

Austenite

Mo

Cr

Mn

Ni

Mo

Cr

Mn

Ni

A. M.

BM HAZ WM

4.48 4.02 4.31

26.30 25.25 25.61

0.65 0.74 0.71

6.35 6.63 6.48

3.08 3.22 3.11

23.46 24.24 23.84

0.86 0.77 0.80

8.28 7.78 8.01

1020 ◦ C

BM HAZ WM

4.88 4.24 4.46

27.04 25.36 26.59

0.62 0.73 0.69

5.77 6.51 6.37

3.25 3.19 3.13

23.58 24.08 23.65

0.83 0.78 0.81

8.07 7.81 7.86

1050 ◦ C

BM HAZ WM

4.75 4.33 4.59

26.53 25.59 26.23

0.63 0.72 0.65

5.98 6.48 6.17

3.17 3.14 3.10

23.34 23.59 23.52

0.83 0.72 0.82

8.16 7.94 8.08

1080 ◦ C

BM HAZ WM

4.51 4.41 4.48

26.17 26.09 26.13

0.67 0.68 0.67

6.25 6.35 6.15

3.04 3.05 3.07

23.32 23.47 23.38

0.83 0.85 0.84

8.34 8.25 8.28

1100 ◦ C

BM HAZ WM

4.46 4.35 4.38

26.03 25.79 25.93

0.70 0.74 0.71

6.32 6.37 6.31

2.93 2.94 2.89

23.31 23.32 23.34

0.82 0.85 0.84

8.45 8.67 8.58

1150 ◦ C

BM HAZ WM

4.34 4.22 4.27

25.87 25.37 25.49

0.73 0.75 0.73

6.42 6.40 6.48

2.65 2.83 2.79

23.29 23.31 23.32

0.85 0.87 0.85

8.57 8.72 8.68

AM, As-welded metal.

2.2. Microstructure characterization To observe the optical micrographs of the WM, HAZ, and BM of the weldments, the polished specimens were electrochemically etched in 30 wt% KOH electrolyte, which made the austenite phase light and the ferrite phase dark. The volume fraction of the ferrite and austenite phases was evaluated carefully by means of a CARL ZEISS optical microscope equipped with a KS400 quantitative metallographical analysis system. The final value was the average of at least 12 measurements. To further identify nitrides or carbides, electrolytic etching was also done at 2.5 V for 10 s in 10% oxalic acid solution before metallographic investigations. Nitrides and carbides are visible as mottled dark areas after this etching process [29–31]. Besides SEM measurement, a JEOL JEM2100F transmission electron microscope(TEM) was also used to identify if there were some secondary phases or not. The chemical compositions of the austenite and ferrite phases were measured by using energy dispersive X-ray spectroscopy (EDS) and the pit morphologies were obtained by a scanning electron microscopy (SEM, FEI Quarter 400) with a Robinson backscattered electron detector. To reduce the effects of concentration fluctuations, more than 10 measurements were carried out for a single specimen. 2.3. Electrochemical measurements All measurements were carried out with a PARSTAT 4000 in a three-electrode cell held in a water bath. The counter electrode was a platinum foil, and a saturated calomel electrode (SCE) were used as the reference electrode. All potentials were given against the SCE. The test solution, 1 mol/L NaCl, was made up of analytical grade reagent and distilled water. The working electrodes were mounted in epoxy resin as shown in Fig. 1. To avoid crevice corrosion, the specimen/resin interfaces were sealed with silica gel sealant and dried in air [32–34]. Potentiodynamic polarization tests were carried out at a scanning rate of 1.33 mV/s from −1500 mV (SCE) to the pitting potential (Epit )SCE at 80 ◦ C. To ensure the reproducibility of the results, experiments were repeated at least three times under the same experimental condition [35,36]. Potentiostatic polarization measurements were also conducted at 850 mV SCE for evaluating the critical pitting temperature (CPT) of the test specimens [37,38]. The heating rate of the electrolyte was kept at

Fig. 3. (a) SEM micrograph of the as-welded HAZ after etching in 10% oxalic acid solution, (b) TEM micrograph of the as-welded HAZ.

1 ◦ C/min. The electrolyte temperature was defined as the CPT of the test specimens when the current density arrived to 100 ␮A/cm2 . The electrolyte was bubbled with pure nitrogen gas to get rid of the dissolved oxygen gas throughout the test. The test results must be eliminated if any crevice corrosion has been observed on the

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Fig. 4. Optical micrographs of the test specimens heat-treated at: (a) 1020 ◦ C, (b) 1080 ◦ C (c) 1150 ◦ C for 3 min after etching in 30 wt% KOH electrolyte.

Fig. 5. Evolution of the ferrite content as function of heat treatment temperature (A. M., As-welded metal, HT, Heat-treated).

surface of the test specimens. All the CPT values were an average of at least three measurements for the same specimen.

3. Results and discussion 3.1. Microstructure of as-welded metal Fig. 2(a–c) illustrates the optical micrographs of the different zones of the as-welded metal. The typical ferrite-austenite duplex

structure was present in the BM. The austenite phase was depicted as the white region, whereas the gray region represented the ferrite phase. Both the ferrite and austenite phases elongated along the rolling direction together with a banded structure. The ferrite matrix in the HAZ and WM were decorated with grain boundary austenite (GBA) that formed at the prior-ferrite grain boundaries and intragranular austenite (IGA) that precipitated in the ferrite grains, whereas the Widmanstätten austenite (WA) grew into the grains from the GBA. The morphology of the austenite phase in

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Fig. 6. Phase diagram of UNS S32750 SDDS calculated by Thermo-Calc software (version N, TcFe 6 database).

Fig. 7. Polarization curves of the specimens tested in 1 mol/L NaCl electrolyte at 80 ◦ C.

the HAZ and WM was irregular. The GBA was coarser than the WA and IGA, thereby suggesting the instability of the ferrite/austenite interface [39]. Fig. 2 presents a lesser austenite content in the HAZ than in the WM, which is inconsistent with the results reported by Kim et al. [23]. The main chemical compositions of the as-welded

metal are listed in Table 3, which presents that more chromium and molybdenum were partial to the ferrite phase and the austenite phase was enriched in nickel in the BM. Conversely, phase stabilizing elements such as chromium, nickel, and molybdenum in the HAZ heavily deviated from the equilibrium state due to the

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Fig. 8. Chrono-amperometry as a function of the electrolyte temperature.

rapid heating and cooling cycles. The significant deviation from the equilibrium state observed in the HAZ may reduce the corrosion resistance of the UNS S32750 SDDS welds. Fig. 3 presents the SEM and TEM micrographs of the as-welded HAZ, wherein no apparent secondary phases were observed to precipitate in the test specimens under the above-mentioned welding parameters.

3.2. Influence of short-time heat treatment 3.2.1. Microstructure evolution and distribution of typical matrix elements Fig. 4 illustrates the microstructure evolution of the test specimens after heat treatment at different temperatures for 3 min. The austenite content in the BM slightly decreased with increasing temperature due to the transformation from the austenite phase into the ferrite phase. A higher quantity and coarser austenite phase was formed in the HAZ and WM as the heat treatment temperature raised to 1080 ◦ C as compared to the BM. The HAZ and WM presented more ferrite phases above 1080 ◦ C. The evolutions of the ferrite volume fraction with temperature for the HAZ, WM, and BM are presented in Fig. 5. The ferrite volume fraction of the as-welded metal agreed indicates that HAZ > WM > BM. The microstructure of the as-welded WM was closer to the equilibrium state than that of the as-welded HAZ. After short-time heat treatment, the ferrite volume fraction of the BM exhibits an increase with the rise in temperature, which is consistent with the results calculated by the Thermo-Calc software (Fig. 6). Higher heat treatment temperatures exhibited greater amounts of the ferrite phase. Contrary to the BM and the equilibrium diagram, the ferrite volume fraction in the HAZ and WM first declined and then rose with an increase in temperature. This indicates a lack of compliance in the microstructure evolution of the heat-treated welds with the phase transition principle displayed in the equilibrium phase diagram. The ferrite volume fraction in the as-welded WM dropped from 58.9% to 50.3% for specimens treated at 1050 ◦ C.Above that temperature, the evolution principle of the ferrite volume fraction in the WM was similar to that in the BM. The minimum ferrite content in the HAZ

was obtained at 1080 ◦ C, which indicates the WM microstructure arrived at the equilibrium state in advance as compared to the HAZ. Table 3 presents the average contents of important alloying elements such as chromium, molybdenum, nickel, and manganese in the ferrite and austenite phases for specimens heated at different temperatures. It is clear that the chromium and molybdenum contents in the ferrite phase of the BM were diluted due to the expanding ferrite phase [6,24,29]. On the other hand, the chromium and molybdenum contents of the austenite phase in the BM decreased with an increase in temperature given the large consumption of chromium and molybdenum in the matrix from the expansion of the ferrite phase. However, chromium exhibited a slower droop rate than molybdenum due to its strong affinity to nitrogen enriched in the austenite phase [6]. The nickel and manganese contents in the austenite phase increased with the increase of temperature because more ferrite-stabilizers were removed during the transformation from austenite to ferrite. As for the aswelded HAZ and WM, the partitioning of alloying elements in either the ferrite or austenite were heavily disturbed. As shown in the previous literature [24,29], the ferrite- or austenite-stabilizing elements did not have enough time to return to their own phases during welding. The alloying elements of chromium, nickel, and molybdenum were nearly equally distributed between the ferrite and austenite phases in the HAZ. After short-time heat treatment, the ferrite- or austenite-stabilizing elements in the HAZ and WM gradually returned to their own phases, which presented an obvious phase transformation. The HAZ reached the equilibrium state at about 1080 ◦ C, whereas the WM reached the equilibrium state at about 1050 ◦ C. Above 1080/1050 ◦ C, the evolution principle of the alloying elements in the HAZ/WM was similar to that in the BM, which was compatible with the microstructure analysis results. 3.2.2. Pitting corrosion resistance The potentiodynamic anodic polarization technique was used to evaluate the effects of the short-time heat treatment temperature on the pitting corrosion resistance of the test specimens. The anodic polarization curves for the test specimens are presented in Fig. 7. The similarity of anodic polarization curves indicates

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Table 4 Electrochemical parameters of the test specimens after three times potentiodynamic polarization measurements in 1 mol/L NaCl solution at 80 ◦ C. Sample

Ecorr (V vs SCE)

Epit (V vs SCE)

Epit − Ecorr

AM 1020 ◦ C 1050 ◦ C 1080 ◦ C 1100 ◦ C 1150 ◦ C

−0.615 ± 0.010 −0.707 ± 0.006 −0.741 ± 0.016 −0.783 ± 0.020 −0.810 ± 0.007 −0.704 ± 0.016

0.204 ± 0.007 0.181 ± 0.017 0.237 ± 0.028 0.287 ± 0.031 0.249 ± 0.014 0.163 ± 0.042

0.819 0.888 0.978 1.070 1.059 0.867

AM, As-welded metal.

an initial activation in all the specimens followed by inactivation until a pitting potential, (Epit )SCE , was obtained [40,41]. The passive current density of the as-welded specimen was higher than that of the heat-treated ones. The appearance of the current peaks triggered under (Epit )SCE suggests the test specimens were attacked by the metastable pits [42,43]. The results of the polarization measurements are listed in Table 4. The values of (Epi t )SCE and the pit nucleation resistance, (Epit -Ecorr )SCE , exhibited a noticeable improvement after short-time heat treatment as compared to the as-welded specimens. The values of (Epit )SCE and (Epit − Ecorr )SCE first increased and then decreased with the temperature increase. The maximum values of (Epit )SCE and (Epit − Ecorr )SCE were obtained for the specimens heat-treated at 1080 ◦ C, which validates 1080 ◦ C as the appropriate 3 min heat treatment temperature for the UNS S32750 SDDS welds to ensure the highest corrosion resistance. Although the potentiodynamic polarization test effectively evaluates the effects of the short-time heat treatment temperature on the pitting corrosion resistance of the test specimens, it is difficult to accurately measure the CPT. Therefore, the potentiostatic critical pitting temperature technique was also executed to obtain an accurate CPT given its good reproducibility, sensitivity, and efficiency [44,45]. Fig. 8 presents the typical curve plots of the current density versus the electrolyte temperature. A current density lower than 1 × 10−6 A/cm2 was exhibited during the initial heating stage, which suggests the presence of passive film protection on the specimen surfaces. The current density sharply ascended near the CPT, which indicates the existence of stable pits. Table 5 lists the CPT value of each specimen. The CPT increased from 67 ◦ C to 74 ◦ C as the heat treatment temperature increased from 1020 ◦ C to 1080 ◦ C, though it decreased following the succeeding increase in temperature. The highest CPT value for the welded alloy was obtained following heat treatment at 1080 ◦ C, which was consistent with the data measured by the potentiodynamic anodic polarization technique. 3.2.3. Pit morphology and pitting resistance of single phase Fig. 9 presents the pit morphologies on the test specimen surfaces after CPT tests. Corrosion pits were found to occur in the ferrite phase in the as-welded HAZ, which deems the ferrite phase in the as-welded HAZ as the weak phase. As shown in our literature [29–32], this was explained by the difference in the pitting resistance equivalent numbers (PRENs) between the ferrite and austenite phases. A higher PREN implies a higher pitting corrosion resistance. All the PRENs were uniformly calculated as PREN = %Cr + 3.3%Mo + 16%N − Mn% [6,7]. By combining the PREN equation and the data listed in Table 3, the PREN values of the single phase were obtained. The nitrogen content in the ferrite phase was assumed to have a saturation value of 0.05%, whereas the rest were partitioned to the austenite phase due to the insensitivity of the energy-dispersive X-ray analysis techniques to nitrogen [7,8]. The calculated PREN value of the single phase versus heat treatment temperature is presented in Fig. 10. The pitting corrosion resistance of the austenite phase in the BM was lower than that of the ferrite phase due to the wide difference in chromium and molybdenum contents between the two phases. Opposite results were observed

Fig. 9. SEM morphologies of metastable pit (a) and (b) stable pit on the surface of the as-welded metal after CPT tests in 1 mol/L NaCl electrolyte.

in the HAZ and WM because of the redistribution of the alloying elements between the ferrite and austenite phases. On one hand, the chromium, nickel, and molybdenum alloying elements in the HAZ were nearly equally distributed between the ferrite and austenite phases. On the other hand, the nitrogen content of the austenite was higher than that of the ferrite given that nitrogen as one of the substitutional elements was easily partitioned into the austenite phase [46]. Therefore, the ferrite phase in the HAZ exhibited the lowest pitting corrosion resistance, and more metastable pits were observed to form in the ferrite phase in the as-welded HAZ. Fig. 11 illustrates the SEM morphologies of the metastable pits on the test specimen surfaces after CPT tests. Although all the corrosion pits were formed in the HAZ, the preferential nucleation sites were different for specimens heat-treated at different temperatures. Most of the metastable pits were located in the ferrite phase for the specimens heat-treated below 1080 ◦ C, whereas corrosion pits mainly initiated at the ferrite/austenite interface for the specimens heat-treated at 1080 ◦ C. Above that temperature, the corrosion pits returned to the ferrite phase. This results could be explained by that while the pitting corrosion resistance of the

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Table 5 CPT of test specimens after six times potentiostatic polarization measurements in1 mol/L NaCl electrolyte. Sample ◦

CPT ( C)

AM

1020 ◦ C

1050 ◦ C

1080 ◦ C

1100 ◦ C

1150 ◦ C

65 ± 0.69

68 ± 0.41

71 ± 0.35

74 ± 0.21

72 ± 0.43

69 ± 0.58

Fig. 10. PREN values of single phase as a function of heat treatment temperature. (a) BM, (b) WM, and (c) HAZ.

ferrite phase in the HAZ gradually increased with the increase in temperature following the return of chromium and molybdenum, the partitioning of the alloying elements in the HAZ still deviated from the equilibrium state for heat treatment temperatures lower than 1080 ◦ C. Therefore, the ferrite phase in the HAZ was still the weakest phase for the specimens heat-treated with above mentioned conditions. As presented in Table 3, an increase in the diffusion rate of the alloying elements allowed the HAZ of the test specimens heat-treated at 1080 ◦ C for 3 min to nearly arrive at the equilibrium state. The ferrite phase in the HAZ exhibited the highest pitting corrosion resistance at 1080 ◦ C. Furthermore, both phases of the HAZ exhibited a similar pitting corrosion resistance at this heat treatment temperature as shown in Fig. 10. Above 1080 ◦ C, the chromium and molybdenum contents in the ferrite phase became diluted due to the expanding ferrite phase and the ferrite pitting corrosion resistance in the HAZ once again decreased. Therefore, specimens heat-treated near 1080 ◦ C for 3 min are predicted to possess the highest pitting corrosion resistance according to the cask effect. Furthermore, almost all corrosion pits readily initiated on the oxide mixed inclusions of the weak phase. As shown in Table 6, these inclusions were rich in elements such as oxygen, aluminum,

Table 6 Mean chemical compositions of the test specimens (wt%) after ten times EDS measurements. Element

O

Al

Si

Ca

Wt%

13.91 ± 0.44

30.44 ± 0.835

11.46 ± 0.01

43.03 ± 0.525

silicon, and calcium. This results are consistent with previous literature [29,47,48]. 4. Conclusions The effects of the short-time heat treatment on the microstructure evolution and pitting corrosion behavior of UNS S32750 SDSS welds were investigated by morphological observation and electrochemical detection. The significant deviation from the equilibrium state observed in the HAZ reduced the corrosion resistance of the UNS S32750 SDSS welds. The ferrite phase in the as-welded HAZ was the weak phase for the corrosion pits to attack. After shorttime heat treatment, the phase-stabilizing elements in the UNS S32750 SDSS welds gradually returned to their own phases, and the

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Program of Shanghai Education Commission (grant No. 14YZ090), and Shanghai University of Engineering Science Innovation Fund (15KY0515). References

Fig. 11. SEM morphologies of metastable pit on the surface of the specimens heattreated at 1020 ◦ C(a), 1080 ◦ C(b) and 1150 ◦ C(c) for 3 min after CPT tests in 1 mol/L NaCl electrolyte.

ferrite content in the HAZ and WM presented an obvious decline until the HAZ and WM reached their equilibrium states. The HAZ and WM reached their equilibrium states at 1080 ◦ C and 1050 ◦ C, respectively. The corrosion pits readily nucleated on the oxide mixed inclusions of the weak phase. The highest pitting corrosion resistance of the UNS S32750 SDSS welds was obtained after heat treatments at 1080 ◦ C for 3 min. Acknowledgements The authors would like to thank Minjie Yan for his help with SEM/EDX analysis. This work is supported by National Natural Science Foundation of China (grant no. 71401106), Shanghai Natural Science Foundation (grant no. 14ZR1418700), Shanghai First-class Academic Discipline Project (grant no. S1201YLXK), and Innovation

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