Intergranular corrosion behavior associated with delta-ferrite transformation of Ti-modified Super304H austenitic stainless steel

Intergranular corrosion behavior associated with delta-ferrite transformation of Ti-modified Super304H austenitic stainless steel

Corrosion Science 90 (2015) 347–358 Contents lists available at ScienceDirect Corrosion Science journal homepage: www.elsevier.com/locate/corsci In...

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Corrosion Science 90 (2015) 347–358

Contents lists available at ScienceDirect

Corrosion Science journal homepage: www.elsevier.com/locate/corsci

Intergranular corrosion behavior associated with delta-ferrite transformation of Ti-modified Super304H austenitic stainless steel Guanshun Bai, Shanping Lu ⇑, Dianzhong Li, Yiyi Li Shenyang National Laboratory for Materials Science (SYNL), Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China

a r t i c l e

i n f o

Article history: Received 11 March 2014 Accepted 23 October 2014 Available online 29 October 2014 Keywords: A. Stainless steel B. Electrochemical calculation B. SEM B. TEM C. Intergranular corrosion

a b s t r a c t A double loop electrochemical potentiokinetic reactivation (DL-EPR) test was conducted to investigate the relationship between the evolution of delta-ferrite and the intergranular corrosion (IGC) of Ti-modified Super304H, which was aged at 650 °C for 4–500 h. Scanning electron microscopy and transmission electron microscopy were adopted to analyze the evolution of delta-ferrite. The results indicated that a higher fraction of delta-ferrite with poor stability increased the IGC sensitisation of Ti-modified Super304H. Moreover, the self-healing of the sensitisation of Ti-modified Super304H occurred after 48 h due to the diffusion of chromium atoms mainly from the adjacent primary austenite rather than the delta-ferrite. Ó 2014 Elsevier Ltd. All rights reserved.

1. Introduction Austenitic stainless steels are widely used as a structural material in steam generating plants as piping and superheating tube materials due to their good combination of mechanical, corrosion resistance and fabrication properties at a relatively low cost [1–3]. However, the intergranular corrosion (IGC) and intergranular stress corrosion cracking (IGSCC) of austenitic stainless steels exposed to the temperature range of 550–800 °C have always been a conventional problem [4–6]. These corrosion failures are attributed to the formation of Cr-depleted zones adjacent to Cr-rich M23C6 carbides and/or other Cr-rich compounds, such as the sigma phase and chi phase at grain boundaries [5,7–12]. The methods to determine the susceptibility to IGC of austenitic stainless steels include conventional tests defined by ASTM A-262 and electrochemical techniques, such as the double loop electrochemical potentiokinetic reactivation (DL-EPR) [13–15], dynamic electrochemical impedance spectroscopy (DEIS) technique [16,17], atomic force microscopy (AFM)–local impedance spectroscopy (LIS) [18] and electrochemical noise (EN) technique [19]. The conventional tests for IGC are qualitative, destructive and restrained under certain terms in their application. For example, the ferric sulfate-sulfuric acid test and copper–copper sulfate16% sulfuric acid test are not suitable for detecting the susceptibility to IGC associated with the sigma phase [20]. Compared with the conventional tests, the electrochemical techniques to evaluate the ⇑ Corresponding author. Tel.: +86 24 23971429; fax: +86 24 83970095. E-mail address: [email protected] (S. Lu). http://dx.doi.org/10.1016/j.corsci.2014.10.031 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.

sensitivity to IGC of austenitic stainless steels are rapid, quantitative and non-destructive [2,13]. At equilibrium, austenitic stainless steels consist of a single austenite phase at room temperature. However, delta-ferrite always forms as a component of the microstructure in austenitic stainless steel due to non-equilibrium solidification and the fabrication methods, including heat treatment and welding. Extensive research has been conducted on delta-ferrite in austenitic stainless steels in terms of the delta-ferrite number prediction methods [21–23], the effects of delta-ferrite on mechanical properties [24] and the hot-workability, such as the susceptibility to hot cracking during welding [25,26], and the influences of delta-ferrite on corrosion resistance properties [27–29]. These studies yielded the following conclusions. Small amounts of delta-ferrite are desirable to prevent hot cracking in the welded austenitic stainless steels and to decrease the crack growth rate in stress corrosion cracking [25–27]. Unfortunately, delta-ferrite in austenitic stainless steel can worsen fatigue properties [30], result in a loss of impact toughness and tensile ductility due to the transformation of delta-ferrite to sigma phase [31] and reduce the pitting corrosion resistance [29]. According to the previous works [28,32,33], the delta-ferrite in austenitic stainless steel decomposes into a Cr-rich phase and secondary austenite phase during the aging process. Therefore, the influence of the delta-ferrite transformation on the intergranular corrosion behavior of austenitic stainless steel needs to be investigated to understand the corrosion resistance of austenitic stainless steel comprehensively. To date, several studies [28,34,35] on the effect of delta-ferrite on susceptibility to IGC of austenitic stainless steels have been published. However, the techniques to determine

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Table 1 Chemical composition of Ti-modified Super304H austenitic stainless steel (wt.%). C

Cr

Ni

Cu

Nb

N

Si

Mn

Mo

Ti

Sample

Average grain size/lm

Delta-ferrite fraction/vol.%

0.10

18.79

9.75

3.24

0.49

0.042

0.26

0.95

0.36

0.58

S1 S2 S3

30.3 ± 12.8 25.2 ± 11.6 31.0 ± 14.0

6.1 ± 0.7 3.9 ± 0.9 0.9 ± 0.4

Table 2 Heat treatment conditions and the corresponding designation of specimens of Ti-modified Super304H used in this study.

a b

Table 3 The average grain size and delta-ferrite fraction of S1, S2 and S3.

Heat treatment condition

Designation

1250 °C  1.5 h(WC)a 1250 °C  1.5 h(WC) + 1200 °C  2 h(WC) + cold rolling (40%)b + 1150 °C  1 h(WC) 1250 °C  1.5 h(WC) + 1200 °C  13 h(WC) + cold rolling (55%)b + 1150 °C  1 h(WC)

S1 S2 S3

WC: water cooling. X%: reduction in the thickness direction.

IGC in these studies are qualitative and insensitive to evaluate the IGC associated with the sigma phase, and data on using the electrochemical techniques to determine the effect of delta-ferrite on intergranular corrosion of austenitic stainless steel quantitatively are lacking. In addition, works [34] on the self-healing (desensitisation) of IGC in austenitic stainless steel containing delta-ferrite are rare, and the self-healing behavior of austenitic stainless steel

associated with the delta-ferrite transformation during the aging process has not been studied systematically. The austenitic stainless steel Super304H was introduced by Sumitomo Metal Industries of Japan and has been widely used for superheater and reheater materials in ultra-supercritical fossil fuel fired boilers due to its superior creep rupture strength and resistance to high temperature oxidation and corrosion [36–38]. The mechanical properties of Super304H have been extensively investigated [38–43]. However, the data available on the corrosion behavior is limited, especially with regard to the susceptibility to IGC of Super304H [44]. In this study, the IGC of Ti-modified Super304H was investigated using a DL-EPR test. Both scanning electron microscopy (SEM) and transmission electron microscopy (TEM) with energy dispersive spectroscopy (EDS) were adopted to analyze the evolution of the delta-ferrite and precipitates. The relationship between the delta-ferrite transformation and IGC behavior (both sensitisation and desensitization) was clarified in the Ti-modified austenitic

Fig. 1. Micrographs of (a) S1, (c) S2 and (e) S3. Grain size distribution and average grain size of (b) S1, (d) S2 and (f) S3.

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Fig. 2. Scanning electron micrographs of specimens containing 6.1% delta-ferrite (S1) aged at 650 °C for: (a) solution annealed, (b) 4 h, (c) 12 h, (d) 48 h, (e) 100 h and (f) 500 h. The inset of each picture shows the austenite/austenite grain boundary segment.

stainless steel Super304H aged at 650 °C for up to 500 h based the analysis of the change in the DL-EPR test data due to the microstructural evolution of the delta-ferrite phase during the aging process.

2. Experimental 2.1. Materials and heat treatments The chemical composition of Ti-modified Super304H austenitic stainless steel used in this research is given in Table 1. The steel was produced by vacuum-melting and then hot forged into 40 mm  40 mm square bars in the temperature range 1200– 900 °C. Subsequently, the bars were solution annealed at 1250 °C for 1.5 h and then rapid water quenched to prevent carbide precipitation. Diffusion annealing at a high temperature with different holding times and cold rolling with different strains were adopted to generate samples with different fractions of delta-ferrite but a similar average grain size and grain size distributions. The heat treatment conditions and the corresponding designation of each specimen are given in Table 2. The samples were etched at 5 V in

a 60% HNO3 solution to measure the grain sizes. They were etched at 10 V in a 20% NaOH solution to measure delta-ferrite fraction. To ensure that the measurement of the delta-ferrite fraction was reproducible, twenty micrographs of each sample were analyzed and the average fraction was considered. The samples with different fractions of delta-ferrite were then aged at 650 °C for 0, 4, 12, 24, 48, 100 and 500 h. The microstructure of each sample was etched at 5 V in a 20% NaOH solution for 5 s to observe the change in delta-ferrite. The samples were etched at 5 V in a 10% oxalic acid solution for 45 s to observe austenite grain boundary. The etched samples were examined via optical microscopy (OM) and scanning electron microscopy (SEM). The precipitates in the aged samples were identified using transmission electron microscopy (TEM) with energy dispersive spectroscopy (EDS). 2.2. Intergranular corrosion tests To quickly and qualitatively identify the IGC attack, the aged specimens were etched at 1 A/cm2 for 90 s in 10% oxalic acid solutions according to the ASTM A262 practice A standard [20]. DL-EPR tests were conducted to examine the intergranular corrosion resistance of the samples containing different fractions of

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Fig. 3. Scanning transmission electron microscopy (STEM) bright-field images of (a) decomposition of delta-ferrite and (b) c/c grain boundary in the specimen containing 6.1% delta-ferrite aged at 650 °C for 48 h, (c) TEM image of chromium carbide at d/c grain boundary, and (d) line profile analysis for Cr at c/c grain boundary (indicated in (b)).

Table 4 Composition of the austenite, delta-ferrite and sigma phases measured by SEM/EDS for the specimen containing 6.1% delta-ferrite (S1) aged at 650 °C for 48 h (wt.%).

Austenite (c) Delta-ferrite (d) Sigma phase (r)

Fe

Cr

Ni

Mo

66.52 65.03 60.90

21.02 28.40 30.29

9.09 4.41 3.95

0.42 0.61 1.86

delta-ferrite using a potentiostat workstation (CHI 760E) with three electrodes according to the standard ISO12732-2006 [14]. A platinum foil and a saturated calomel electrode (SCE) were used as the auxiliary and reference electrodes, respectively. The samples containing different fractions of delta-ferrite were embedded in epoxy resin with an exposure area of 1 cm2. They were successively grinded to 2000 grit using emery papers. During the DL-EPR tests, the working electrode was first immersed in the electrolyte containing 0.5 M H2SO4 + 0.01 M KSCN and cathodically polarized at 0.6 V vs. SCE for 2 min to improve reproducibility. Second, the working electrode was anodically polarized from 0.6 V vs. SCE to +0.3 V vs. SCE when the steady-state open circuit potential (e.g., approximately 0.4 V vs. SCE) was reached (approximately 10–20 min) and then cathodically polarized to 0.6 V vs. SCE from +0.3 V vs. SCE at a scan velocity of 6 V/h. A water bath was used to maintain the electrolyte temperature of 25 °C. At least three separate samples were tested for each group to ensure the reproducibility of the DL-EPR test results. After the DL-EPR tests, the optical microscopy (OM) and scanning electron microscopy (SEM) were used to observe the morphology of sample. The susceptibility to intergranular corrosion can be evaluated by the reactivation ratio (Rr), which is defined as follows:

ir  100% ia Q Rr ¼ r  100% Qa

Rr ¼

ð1Þ ð2Þ

where ir is the reactivation current density peak (the maximum current density of the reactivation scan), ia is the activation current density peak (the maximum current density of the reactivation scan) and, Q r and Q a are the charge for the reactivation scan and the charge for the activation scan in the DL-EPR curve, respectively. 3. Results and discussions 3.1. Microstructure characterisations Fig. 1 shows the micrographs, average grain size and the grain size distributions of the samples containing different fractions of delta-ferrite prior to the aging heat treatment. In the micrographs of each sample, the dark phase indicates delta-ferrite (d) and the bright phase is austenite (c). The average grain size and delta-ferrite volume fraction of the samples (S1, S2 and S3) are summarised in Table 3. Fig. 2 presents the microstructures of specimens containing 6.1% delta-ferrite aged at 650 °C for various times. The inset in Fig. 2 shows the austenite/austenite (c/c) grain boundary segment. The SEM micrograph indicates that the solution-treated sample contained a darker delta-ferrite phase and lighter austenite phase and no precipitates at grain boundaries according to the SEM/EDS analysis, as shown in Fig. 2(a). However, a small number of sigma (r) phases were formed at the delta-ferrite/austenite (d/c) grain boundaries, while no visible precipitates appeared at the c/c grain boundaries after aging at 650 °C for 4 h (as shown in

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Fig. 4. Microstructure after a 10% oxalic acid etch test on the samples containing 6.1% delta-ferrite (S1) aged at 650 °C for different times: (a) solution treated, (b) 4 h, (c) 24 h, (d) 48 h, (e) 100 h and (f) 500 h.

Fig. 2(b)). This structure is attributed to the higher diffusion rate of Cr atoms in delta-ferrite than in austenite [45]. The appearance of the lighter secondary austenite phase (c2) inside the darker deltaferrite indicated the decomposition of the delta-ferrite aged at 650 °C. This result was later confirmed by the TEM analysis, as shown in Fig. 3(a). As aging continued (from 4 h to 48 h), the size of the sigma phases precipitated at the d/c grain boundaries increased significantly and several small sigma phases appeared in the delta-ferrite interiors. Moreover, a number of M23C6 carbides appeared at the c/c grain boundaries, as shown in Fig. 2(d). The main composition of the austenite, delta-ferrite and sigma phases identified by SEM/EDS for the specimen aged at 650 °C for 48 h is given in Table 4. When the specimen was aged at 650 °C for 100 h, the deltaferrite completely transformed into sigma phase and the secondary austenite phase (Fig. 2(e)). The sigma phases that appeared inside of the delta-ferrite grew. In addition, the M23C6 carbides at the c/c grain boundaries also grew. As the aging time increased to 500 h (Fig. 2(f)), the number of sigma phases inside the delta-ferrite decreased but they grew in size due to the Ostwald ripening of sigma phases. Moreover, the M23C6 carbides at the c/c grain boundaries coarsened. The decomposition of delta-ferrite and the crystalline structure of the intermetallic phases were determined by TEM. Fig. 3 shows a TEM micrograph of the decomposition of the delta-ferrite in the sample containing 6.1% delta-ferrite aged at 650 °C for 48 h and the electron diffraction patterns of each phase. Fig. 3(a) shows the scanning transmission electron microscopy (STEM) bright-field

image of the decomposition of the delta-ferrite. The phases decomposed from delta-ferrite were identified as a secondary austenite phase and sigma phase. Notably, M23C6 also precipitated at the d/c grain boundary, as shown in Fig. 3(c). Both the M23C6 and Cr-depleted zones can be observed along the c/c grain boundary, as indicated in Fig. 3(b) and (d). In particular, the line profile over a straight line in Fig. 3(b) demonstrates that the Cr-depleted zones are asymmetric in the vicinity of M23C6 for the sample aged at 650 °C for 48 h, as shown in Fig. 3(d). This result is similar to the results reported by Kaneko et al. [11] and Sourmail et al. [46]. 3.2. Oxalic acid etch test Fig. 4 shows the microstructures of the samples containing 6.1% delta-ferrite aged at 650 °C for various times etched using a 10% oxalic acid solution. The microstructure of the solution-treated specimen is shown in Fig. 4(a). According to the ASTM A262-A standard [20], this figure shows a step structure. The d/c and c/c grain boundaries only show some steps, and no ditches were observed. Fig. 4(b) demonstrates the microstructure of the specimen aged for 4 h. Both the d/c and c/c grain boundaries showed some ditches in addition to steps, but did not show any single grains completely surrounded by ditches. As the aging time increased (from 4 h to 48 h), the number of ditches at the c/c grain boundaries increased significantly, and some grains were completely surrounded by ditches as shown in Fig. 4(b)–(d). In particular, the delta-ferrite was corroded severely to result in a ditch

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Fig. 5. Typical DL-EPR curves of Ti-modified Super304H austenitic stainless steel containing different fractions of delta-ferrite (S1, S2 and S3) aged at 650 °C for various times: (a) solution treated, (b) 4 h, (c) 12 h, (d) 48 h, (e) 100 h and (f) 500 h.

structure. The degree of sensitisation of the specimen appeared to be positively correlated with the aging time. However, when the sample was aged 100 h, only a limited number of ditches could be observed at the c/c grain boundaries in addition to the steps, which indicated that desensitisation had occurred. The specimen aged for 500 h (in Fig. 4(f)) showed only steps, and few ditches could be observed at the c/c grain boundaries, which indicated that prolonging the aging time from 100 h to 500 h decreased the degree of sensitisation of the specimen. 3.3. DL-EPR test Fig. 5 shows the typical DL-EPR curves of Ti-modified Super304H austenitic stainless steels containing different fractions of deltaferrite (S1, S2 and S3) aged at 650 °C for various times. All specimens showed a wide passivity range from 0.15 V vs. SCE to +0.3 V vs. SCE. When the S1, S2 and S3 were aged at 650 °C for the same time, the activation current density peak (ia) was almost independent of the fraction of delta-ferrite, while the magnitude of the reactivation current density peak (ir) varied with the fraction of delta-ferrite. The reactivation current density peak (ir) in the reactivation scan was attributed to the chromium-depleted regions

induced by the precipitation of the sigma phase and chromium carbide M23C6. A reactivation current density peak (ir) was not apparent in the reactivation scan of the solution-treated samples (Fig. 5(a)), which indicates a lack of a chromium-depleted zone in the solution-treated samples. After being aged at 650 °C as shown in Fig. 5(b)–(e), the reactivation current density peaks in the DL-EPR curves indicate chromium-depleted zones at grain boundaries in the aged samples. In addition, ir and ia did not occur at the same potential, which was attributed to the ohmic resistance drop [47]. The difference between the potentials of ir and ia was also reported in earlier studies of duplex stainless steel [48–50]. A second anodic peak (i2) was observed in the passive range in the activation scan of DL-EPR curves for all specimens at different heat treatment conditions, as shown in Fig. 5(a)–(f). According to Ref. [51], the origin of the second anodic peak was attributed to the dissolution of the copper-rich corrosion product layer in sulfuric acid, and the chromium-depleted zone at the grain boundary had a negligible effect on the second anodic current peak. Therefore, the second anodic current peak did not influence the DL-EPR test of the IGC, which was induced by the chromiumdepleted zone around the grain boundary of Ti-modified Super304H austenitic stainless steel in this study.

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Fig. 6. (a) Anodic polarization curve of the samples containing 6.1% delta-ferrite in the 0.5 M H2SO4 + 0.01 M KSCN solution at 25 °C and (b) current transient behavior in the potentiostatic etching tests, at an applied potential of +0.1 V vs. SCE for 300 s in the 0.5 M H2SO4 + 0.01 M KSCN solution at 25 °C, of the samples containing 6.1% delta-ferrite with and without anodic polarization.

In order to better understand the origin of the second anodic peak, the electrochemical tests were performed on the solution treated samples containing 6.1% delta-ferrite in the 0.5 M H2SO4 + 0.01 M KSCN solution at 25 °C. The electrochemical tests included following cases: (a) The samples were anodically polarized from 0.6 V vs. SCE to 0 V vs. SCE at a scan velocity of 6 V/h, and subsequently potentiostatic etched at +0.1 V vs. SCE for 300 s. (b) The samples without anodic polarization were potentiostatic etched at +0.1 V vs. SCE for 300 s. Fig. 6(a) shows the anodic polarization curve of the solution treated sample containing 6.1% delta-ferrite in the 0.5 M H2SO4 + 0.01 M KSCN solution at 25 °C. Fig. 6(b) presents the current transient behavior in the potentiostatic etching tests of the samples containing 6.1% delta-ferrite with and without anodic polarization. A current peak (indicated by arrow in Fig. 6(b)) existed in the current transient curve of the sample with anodic polarization, indicating that the active dissolution occurred at +0.1 V vs. SCE. The current transient curve of the sample without anodic polarization reveals that the austenite and delta-ferrite matrix in the solution treated sample were passivated at +0.1 V vs. SCE as shown in Fig. 6(b), and hence did not contribute to the second anodic peak. The OM image of the sample after anodic polarization from 0.6 V vs. SCE to 0 V vs. SCE at a scan velocity of 6 V/h was shown as Fig. 7(a). It can be seen that a corrosion product layer formed on the surface of the sample after anodic polarization, as shown

353

in Fig. 7(b). The chemical composition of the corrosion product was analyzed using scanning electron microscopy (SEM) with energy dispersive spectroscopy (EDS). According to the results of EDS analysis in Fig. 7(f), the corrosion product layer after anodic polarization was rich in Cu compared with the austenite matrix, indicating that the anodic polarization from 0.6 V vs. SCE to 0 V vs. SCE at a scan velocity of 6 V/h in the DL-EPR test resulted in a Cu-rich corrosion product layer on the surface of the sample. After potentiostatic etching at +0.1 V vs. SCE for 300 s, the corrosion product layer on the sample with anodic polarization was partially dissolved, as shown in Fig. 7(c). The result can be confirmed by the EDS analysis that some regions were still covered by the Cu-rich corrosion product (Fig. 7(d)) while the other regions were not (Fig. 7(e)). That was the reason why a peak occurring in the current transient curve of the sample with anodic polarization in the potentiostatic etching test, as shown in Fig. 6(b). The results indicated that the second anodic peak around +0.1 V vs. SCE in the activation scan of DL-EPR curve was attributed to the dissolution of the copper-rich corrosion product layer. According to the literature [52,53], when more than one current peak presented in the activation and/or reactivation scan of the DL-EPR curve, the charge ratio (Qr/Qa) is a more suitable parameter to determine the degree of sensitisation, which can avoid extra work to separate the peaks related to the dissolution around different deleterious phases. Therefore, the degree of sensitisation, the reactivation ratio (Rr), was calculated by using the charge ratio (Qr/Qa) in this study, as shown in Eq. (2). The values of the reactivation ratio (Rr) measured from the DL-EPR curves of Ti-modified Super304H austenitic stainless steels are summarised in Table 5. The results show that the DL-EPR tests for the evaluation of the IGC of Ti-modified Super304H austenitic stainless steel were reproducible. Fig. 8 shows the comparison of the IGC morphologies observed after the DL-EPR test for the samples containing different fractions of delta-ferrite aged at 650 °C for various aging times. To acquire the details of IGC, the surface of the specimens containing 6.1% delta-ferrite were analyzed by SEM after the DL-EPR test, as indicated in Fig. 9. The solution-treated specimen did not show IGC attack at the c/c grain boundaries (Fig. 9(a)). When the sample was aged for 4 h, IGC was not evident at the c/c grain boundaries, while attack on the d/c grain boundaries was observed due to the formation of sigma phases (as shown in Fig. 9(b)). IGC attack was evident in the austenite phase at the d/c grain boundary. When aged for 12 h, the IGC attack occurred in the delta-ferrite at the d/c grain boundary together with a narrow attack in the c/c grain boundary, as shown in Fig. 9(c). The change in the IGC site at the d/c grain boundary will be discussed in detail in Section 3.4. As the aging time increased (from 12 h to 48 h), a wider IGC attack occurred in the primary delta-ferrite phase at the d/c grain boundary, while a wider attack was formed inside the c/c grain boundary, as shown in Fig. 9(d). However, the attack at the c/c grain boundary of the specimen aged for 100 h (Fig. 9(e)) narrowed compared with that observed for the specimen aged for 48 h. Imperceptible IGC attacks occurred inside the c/c grain boundaries of the specimen when the aging time increased to 500 h, as presented in Fig. 9(f). The DL-EPR test result is consistent with the result of the test using 10% oxalic acid. Fig. 10 presents the variation of the degree of sensitisation as a function of the aging time for the three different fractions of deltaferrite samples. The values of Rr increased rapidly for all specimens when the aging time increased from 0 h (solution-treated) to 12 h at 650 °C. This relationship is mainly related to the rapid precipitation of the sigma phases and chromium carbides, as shown in Fig. 2(a)–(c). As the aging time increased (in the range of 12–48 h), the value of Rr increased at a slower rate. This result arises from the decreased growth rate for the sigma phase due to

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Fig. 7. Surface morphologies of the samples containing 6.1% delta-ferrite after electrochemical tests: (a) OM image of the sample after anodic polarization, (b) SEM image of the region A shown in (a), (c) OM image of the sample with anodic polarization after potentiostatic etching, (d) SEM image of region B shown in (c), (e) SEM image of region C shown in (c), and (f) EDS analysis of the rectangular areas in (b), (d) and (e). The chemical composition of each region is the average of five measurements.

Table 5 Values the reactivation ratio (Rr) measured from the DL-EPR curves of Ti-modified Super304H austenitic stainless steels containing different fractions of delta-ferrite (S1, S2 and S3) aged at 650 °C for various times. Aging time (h)

Rr (%) S1

4 12 24 48 100 500 a

S2

S3

Mean

ra

Mean

r

Mean

r

7.280 10.723 11.263 12.737 1.108 0.733

0.544 0.107 0.325 1.399 0.260 0.494

0.013 1.013 2.020 1.650 0.207 0.107

0.023 0.272 0.235 0.297 0.086 0.067

0.003 1.097 1.797 2.083 0.102 0.030

0.006 0.206 0.189 0.880 0.117 0.028

r: standard deviation.

the consumption of Cr atoms. The highest grade of sensitisation occurred at 48 h. However, the values of Rr for the three different fractions of delta-ferrite samples all decreased at a relatively quicker rate as the aging time increased from 48 h to 100 h. The result indicates that sensitisation self-heals in the samples after 48 h. In addition, the value of Rr for the specimen containing a higher fraction of delta-ferrite is higher than that of the specimens containing a lower fraction of delta-ferrite for the same aging conditions (aging time and temperature). For example, the value of Rr for the specimen containing 6.1% delta-ferrite is 12.737%, which is higher than the 1.650% value for the specimen containing 3.9%

delta-ferrite aged at 650 °C for 48 h. The reasons for these results will be discussed later in Section 3.4. Based on the observation of the morphologies of each sample after the DL-EPR test, the critical value of Rr for the IGC of Ti-modified Super304H is approximately 1.108% under the experimental conditions proposed in this study. The time-delta ferrite fraction-sensitisation (TFS) curve plotted according to the DL-EPR test results is shown in Fig. 11. IGC occurs in the sample aged at the regions indicated as dark circles in the figure. As shown in Fig. 11, txs is the starting time for the sample containing x vol.% delta-ferrite to be sensitized, and t xf is the finishing time for the sample containing x vol.% delta-ferrite to be self-healed at 650 °C.   Dtx ¼ t xf  txs is defined as the self-healing time for the sample containing x% delta-ferrite. The curve indicates that the decrease in the delta-ferrite fraction requires a longer time to be sensitized but shorter time to be healed. For example, the sensitisation of the specimen containing 0.9% delta-ferrite takes longer to begin but shorter to heal, which agrees with the OM observation of the IGC morphologies after the DL-EPR test shown in Fig. 8. 3.4. The effect of delta-ferrite on the sensitisation of Ti-modified Super304H austenitic stainless steel In the present study, the IGC experiments that used both the oxalic acid test and DL-EPR test indicate that a higher delta-ferrite

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Fig. 8. IGC morphologies observed after the DL-EPR test for the samples containing different fractions of delta-ferrite aged at 650 °C for various aging times.

fraction increases the susceptibility of Ti-modified Super304H austenitic stainless steel aged at 650 °C to IGC. The detrimental effects of delta-ferrite on the IGC resistance of Ti-modified Super304H austenitic stainless steel in the present work are discussed in detail as follows. The sigma phase [54] and Cr carbide [55] are well known to preferentially precipitate at the d/c grain boundary. Both the sigma phase and chromium carbide are Cr-rich phases, and their precipitation at the grain boundary can result in Cr-depleted zones adjacent to the grain boundary, which are subsequently prone to IGC at certain conditions. According to the literature [54], the precipitation of the sigma phase in the present study is determined by the interface energy of the d/c grain boundary and the chemical driving force of Cr. Therefore, the increase of the delta-ferrite fraction will increase the number of d/c grain boundaries and offer more preferential sites for the precipitation of the Cr-rich phase, which may result in more a severe susceptibility to intergranular corrosion of the d/c grain boundary of Ti-modified Super304H austenitic stainless steel. The delta-ferrite in our research was unstable and prone to transform to the sigma phase and secondary austenite phase when

aged at 650 °C over a short time (4 h). As shown in Fig. 9(b), IGC attack occurs within the austenite phase at the d/c grain boundary at the beginning of aging process at 650 °C (4 h). Only a few deltaferrite likely decomposed, and the Cr atoms diffusing from the remaining delta-ferrite can replenish the Cr-depleted zones on the ferrite side when aged at 650 °C for a short time. However, when aged up to 12 h, IGC occurs within the delta-ferrite at the d/c grain boundary, as indicated in Fig. 9(c). This phenomenon occurs because the amount of the sigma phase decomposed from the delta-ferrite increases at the d/c grain boundary and leads to more severe Cr depletion on the ferrite side around the sigma phase, as demonstrated in Fig. 2(c), while the amount of retained delta-ferrite is insufficient to transport Cr atoms to the Cr-depleted zones. Thus, the IGC attack occurs on the ferrite side at the d/c grain boundary (Fig. 9(c)). The change in the IGC site at the d/c grain boundary with the aging time indicates that the stability of the delta-ferrite during the aging process plays an important role in the intergranular corrosion behavior of Ti-modified Super304H austenitic stainless steel. If the delta-ferrite is stable and difficult to decompose during aging process at a specific temperature, the Cr-depleted zone along

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Fig. 9. SEM morphologies after the DL-EPR test for the samples containing 6.1% delta-ferrite (S1) aged at 650 °C for various aging times: (a) solution treated, (b) 4 h, (c) 12 h, (d) 48 h, (e) 100 h and (f) 500 h.

the d/c grain boundary can receive Cr atoms from the delta-ferrite quickly due to the high diffusion rate of Cr in the ferrite. However, the beneficial effect of the delta-ferrite as a Cr source will be diminished when the delta-ferrite is unstable and decomposed into the Cr-rich phase and secondary austenite easily. Moreover, the sigma phases decomposed from the delta-ferrite can also give rise to the formation of additional Cr-depleted zones. Therefore, the unstable delta-ferrite during the aging process can aggravate the sensitisation of Ti-modified Super304H austenitic stainless steel. In addition, the onset times of the sensitisation of the c/c grain boundary differ among the samples containing different fractions of delta-ferrite aged at 650 °C, as illustrated in Figs. 8 and 11. The delta-ferrite appears to be able to affect the sensitisation behavior of the c/c grain boundary. As shown in Fig. 3(b) and (d), the sensitisation of the c/c grain boundary is attributed to the formation of chromium carbide. The carbon content is well known to be the predominant factor to control the sensitisation during the intergranular corrosion process of austenitic stainless steel [8]. Carbon is the austenite-forming element, and the solubility of carbon in the austenite is higher than that in the ferrite. Therefore, the carbon concentration in the austenite phases will increase with the delta-ferrite fraction in the Ti-modified Super304H austenitic

stainless steel. The high concentration of carbon can not only enhance the precipitation kinetics of chromium carbides but also increase the number of chromium carbides at c/c grain boundary. As a result, the c/c grain boundary is sensitized more rapidly but desensitized more slowly in the higher fraction of the delta-ferrite sample. Based on the analysis above, a higher fraction of poorly stable delta-ferrite in austenitic stainless steel Ti-modified Super304H increases the number of d/c grain boundaries, the carbon content in the austenite matrix and the amount of the sigma phase at 650 °C, which can severely deteriorate the IGC resistance of the austenitic stainless steel Ti-modified Super304H. 3.5. The self-healing mechanism of IGC of Ti-modified Super304H The IGC sensitisation self-healed in Ti-modified Super304H austenitic stainless steel containing delta-ferrite aged at 650 °C after 48 h, as shown in Fig. 8. In austenitic stainless steel, Cr-depleted regions occur along the grain boundary once the Cr atoms diffuse from the matrix to the grain boundary and form Cr-rich phases, such as Cr carbides and sigma phases. If the diffusion rate of Cr from the matrix to the Cr-depleted zones is higher than the depletion rate of Cr atoms

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Fig. 10. Degree of sensitisation (DOS) as a function of aging time in the specimens containing different fractions of delta-ferrite.

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grain boundary and source significantly more Cr atoms from delta-ferrite, which results in a severely Cr-depleted zone adjacent to the sigma phase. This Cr-depleted zone promotes the decomposition rate of delta-ferrite into the sigma phase and secondary austenite phase. In this case, the Cr-depletion rate is higher than the diffusion rate of Cr from the intact ferrite matrix to the Cr-depleted zones. As the aging time increased (after 48 h), due to the growth rate of sigma phase at d/c grain boundary decreases, the diffusion rate of Cr from the austenite matrix to the Cr-depleted zones is higher than the depletion rate of Cr atoms induced by the growth of the sigma phase at the d/c grain boundary. Eventually, the Cr atoms from the austenite matrix rather than the delta-ferrite replenish the Cr-depleted zones, which can be confirmed by the growth of the sigma phase at the d/c grain boundary toward the austenite phase, as shown in the inset of Fig. 9(f). As discussed in Section 3.4, poorly stable delta-ferrite acts as the Cr source, which can induce the self-healing of the IGC, but this role will be diminished when the sample is aged at 650 °C for 48 h. Therefore, it is the Cr atoms from the primary austenite phase that reduce the susceptibility of Ti-modified Super304H to IGC after 48 h. The conclusions in this study were reached based on the IGC tests for an austenitic stainless steel containing little delta-ferrite (less than 6.1%), and future studies should conduct IGC tests on austenitic stainless steel containing higher fractions of deltaferrite.

4. Conclusions

Fig. 11. Time-delta ferrite fraction-sensitisation (TFS) curve based on the DL-EPR test results.

due to the precipitations of the Cr-rich phases, the extent of Cr-depletion will diminish gradually and the self-healing of sensitisation finally occurs [56]. Therefore, self-healing depends on the precipitation rate of the Cr-rich phases, and the self-healing time depends on the extent of the Cr-depleted zone around the Cr-rich carbide. As mentioned previously, the precipitation kinetics of the Cr-rich phases, such as the sigma phase, depends on the chemical driving force and the energy of the interface. In this work, the precipitation of the sigma phase was mainly determined by the diffusion of Cr atoms and the nature of the d/c grain boundary, while the formation of Cr carbide was strongly related to the nature of the d/c or c/c grain boundary as well as the content of the available carbon to form Cr carbide. In the present study, the self-healing of the c/c grain boundary was related to the precipitation behavior of chromium carbide. For the specimen containing 6.1% delta-ferrite, Cr carbide continued to grow at the c/c grain boundary after 48 h (in Fig. 2(c)–(e)), but the susceptibility to the IGC of the specimen decreased (as shown in Fig. 10). Here, the growth rate of Cr carbide decreased due to the limitation of the carbon consumption after 48 h, which results in a lower Cr-depleted rate induced by the growth of Cr carbide than the diffusion rate of the Cr from austenite matrix to the Cr-depleted zones. In this case, the Cr from the near austenite phases will replenish the depleted Cr at the c/c grain boundary. Moreover, the self-healing of IGC at the d/c grain boundary can be mainly explained by the precipitation behavior of the sigma phases at the d/c grain boundary. When aged at 650 °C for a short time (up to 48 h), the sigma phases precipitate quickly at the d/c

(1) The DL-EPR test method can be recommended to determine the influence of delta-ferrite on the susceptibility of Ti-modified Super304H austenitic stainless steel to IGC quantitatively. (2) The IGC sensitisation of Ti-modified Super304H is caused by the precipitation of the sigma phase and M23C6. (3) The delta-ferrite in Ti-modified Super304H austenitic stainless steel can decompose into a sigma phase and a secondary austenite phase rapidly during aging process at 650 °C, which diminishes the role of delta-ferrite as the Cr source. (4) The IGC site changed with the aging time at the austenite/ ferrite grain boundary in Ti-modified Super304H austenitic stainless steel at 650 °C. The stability of delta-ferrite strongly influences the IGC resistance of Ti-modified Super304H. A higher fraction of poorly stable delta-ferrite can increase the IGC sensitisation of Ti-modified Super304H at 650 °C. (5) The susceptibility of Ti-modified Super304H aged at 650 °C to IGC increases positively correlates with the aging time up to 48 h according to DL-EPR. Further prolonging the aging time beyond 48 h decreases the value of Rr due to the diffusion of Cr atoms from the adjacent primary austenite phase to chromium-depleted regions.

Acknowledgements The authors would like to thank Zhenzhen Peng and Jinmin Liu for their help with the TEM observation. This work is supported by the National High Technology Research and Development Program of China (Grant No. 2012AA03A501). References [1] Y. Yamamoto, M.P. Brady, Z.P. Lu, P.J. Maziasz, C.T. Liu, B.A. Pint, K.L. More, H.M. Meyer, E.A. Payzant, Creep-resistant, Al2O3-forming austenitic stainless steels, Science 316 (2007) 433–436.

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G. Bai et al. / Corrosion Science 90 (2015) 347–358

[2] G.H. Aydogdu, M.K. Aydinol, Determination of susceptibility to intergranular corrosion and electrochemical reactivation behaviour of AISI 316L type stainless steel, Corros. Sci. 48 (2006) 3565–3583. [3] J. Xu, X.Q. Wu, E.H. Han, Acoustic emission response of sensitized 304 stainless steel during intergranular corrosion and stress corrosion cracking, Corros. Sci. 73 (2013) 262–273. [4] R. Singh, S.G. Chowdhury, B.R. Kumar, S.K. Das, P.K. De, I. Chattoraj, The importance of grain size relative to grain boundary character on the sensitization of metastable austenitic stainless steel, Scr. Mater. 57 (2007) 185–188. [5] C.S. Tedmon, D.A. Vermilyea, J.H. Rosolowski, Intergranular corrosion of austenitic stainless steel, J. Electrochem. Soc. 118 (1971) 192–202. [6] A. Abou-Elazm, R. Abdel-Karim, I. Elmahallawi, R. Rashad, Correlation between the degree of sensitization and stress corrosion cracking susceptibility of type 304H stainless steel, Corros. Sci. 51 (2009) 203–208. [7] K. Chandra, V. Kain, R. Tewari, Microstructural and electrochemical characterisation of heat-treated 347 stainless steel with different phases, Corros. Sci. 67 (2013) 118–129. [8] A. Pardo, M.C. Merino, A.E. Coy, F. Viejo, M. Carboneras, R. Arrabal, Influence of Ti, C and N concentration on the intergranular corrosion behaviour of AISI 316Ti and 321 stainless steels, Acta Mater. 55 (2007) 2239–2251. [9] T. Devine, The mechanism of sensitization of austenitic stainless steel, Corros. Sci. 30 (1990) 135–151. [10] E.L. Hall, C.L. Briant, Chromium depletion in the vicinity of carbides in sensitized austenitic stainless steels, Metall. Mater. Trans. A 15 (1984) 793– 811. [11] K. Kaneko, T. Fukunaga, K. Yamada, N. Nakada, M. Kikuchi, Z. Saghi, J.S. Barnard, P.A. Midgley, Formation of M23C6-type precipitates and chromiumdepleted zones in austenite stainless steel, Scr. Mater. 65 (2011) 509–512. [12] S.X. Li, Y.N. He, S.R. Yu, P.Y. Zhang, Evaluation of the effect of grain size on chromium carbide precipitation and intergranular corrosion of 316L stainless steel, Corros. Sci. 66 (2013) 211–216. [13] V. Cihal, R. Stefec, On the development of the electrochemical potentiokinetic method, Electrochim. Acta 46 (2001) 3867–3877. [14] ISO12732, Corrosion of metals and alloys. Electrochemical potentiokinetic reactivation measurement using the double loop method (based on Cihal’s Method), 2006. [15] P. Muri, F.V.V. Sousa, K.S. Assis, A.C. Rocha, O.R. Mattos, I.C.P. Margarit-Mattos, Experimental procedures and sensitization diagnostics of AISI304 Steel by double loop electrochemical potentiodynamic reactivation method, Electrochim. Acta 124 (2014) 183–189. [16] A. Arutunow, K. Darowicki, DEIS assessment of AISI 304 stainless steel dissolution process in conditions of intergranular corrosion, Electrochim. Acta 53 (2008) 4387–4395. [17] A. Arutunow, K. Darowicki, DEIS evaluation of the relative effective surface area of AISI 304 stainless steel dissolution process in conditions of intergranular corrosion, Electrochim. Acta 54 (2009) 1034–1041. [18] A. Arutunow, K. Darowicki, M.T. Tobiszewski, Electrical mapping of AISI 304 stainless steel subjected to intergranular corrosion performed by means of AFM–LIS in the contact mode, Corros. Sci. 71 (2013) 37–42. [19] M.G. Pujar, N. Parvathavarthini, R.K. Dayal, S. Thirunavukkarasu, Assessment of intergranular corrosion (IGC) in 316(N) stainless steel using electrochemical noise (EN) technique, Corros. Sci. 51 (2009) 1707–1713. [20] ASTM A262-10, Standard Practices for Detecting Susceptibility to Intergranular Attack in Austenitic Stainless Steels, ASTM, New York, 2010. [21] T. Gill, R. Dayal, J. Gnanamoorthy, Estimation of delta ferrite in austenitic stainless steel weldments by an electrochemical technique, Weld J 12 (1979). [22] D. Qixun, Y. Ruzeng, The calculation of Td and Vd in austenitic steel, Mater. Charact. 38 (1997) 129–133. [23] Y. Cui, C.D. Lundin, Ferrite number as a function of the Larson–Miller parameter for austenitic stainless weld metals after creep testing, Metall. Mater. Trans. A 35A (2004) 3631–3633. [24] D. Read, H. McHenry, P. Steinmeyer, R. Thomas Jr., Metallurgical factors affecting the toughness of 316 L SMA weldments at cryogenic temperatures, Weld J 59 (1980) 104. [25] T. Ogawa, E. Tsunetomi, Hot cracking susceptibility of austenitic stainlesssteels, Weld J 61 (1982) S82–S93. [26] J. Lippold, W. Savage, Solidification of austenitic stainless steel weldments: Part III–the effect of solidification behavior on hot cracking susceptibility, Weld J 61 (1982) 388. [27] T.R. Lucas, The effect of thermal aging and boiling water reactor environment on Type 316L stainless steel welds, in: Department of Nuclear Science and Engineering, Massachusetts Institute of Technology, 2011. [28] T.M. Devine, Mechanism of intergranular corrosion and pitting corrosion of austenitic and duplex 308 stainless-steel, J. Electrochem. Soc. 126 (1979) 374– 385. [29] P. Manning, D. Duquette, W. Savage, Technical note: the effect of retained ferrite on localized corrosion in duplex 304L stainless steel, Weld J 59 (1980) 260–262.

[30] B.S. Rho, H.U. Hong, S.W. Nam, The effect of d-ferrite on fatigue cracks in 304L steels, Int. J. Fatigue 22 (2000) 683–690. [31] C. Tseng, Y. Shen, S. Thompson, M. Mataya, G. Krauss, Fracture and the formation of sigma phase, M23C6, and austenite from delta-ferrite in an AlSl 304L stainless steel, Metall. Mater. Trans. A 25 (1994) 1147–1158. [32] J.J. Smith, R.A. Farrar, Effect of composition on the transformation behavior of duplex 316-weld metal, J. Mater. Sci. 26 (1991) 5025–5036. [33] C.C. Hsieh, W. Wu, Overview of intermetallic sigma (r) phase precipitation in stainless steels, ISRN Metall. 2012 (2012). [34] Y. Nakao, K. Nishimoto, M. Ishizaki, Influence of delta-ferrite on sensitization of the austenitic stainless steel weld metal, Q. J. Jpn. Weld. Soc. (Jpn.) 9 (1991) 97–104. [35] T.M. Devine, Influence of carbon content and ferrite morphology on the sensitization of duplex stainless-steel, Metall. Trans. A 11 (1980) 791–800. [36] Y. Samaragi, N. Otsuka, H. Senba, S. Yamamoto, Properties of a new 18–8 austenitic stainless steel tube (Super 304H) for fossil fired boilers after service exposure with high elevated temperature strength, Sumitomo Search 56 (1994) 34–43. [37] Y. Sawaragi, K. Ogawa, S. Kato, Development of the economical 18–8 austenitic stainless steel (Super 304H) having high elevated temperature strength for fossil fired boilers, Sumitomo Search 48 (1992) 50–58. [38] K. Ogawa, Y. Sawaragi, N. Otsuka, H. Hirata, A. Natori, S. Matsumoto, Mechanical and corrosion properties of high strength 18% Cr austenitic stainless steel weldment for boiler, ISIJ Int. 35 (1995) 1258–1264. [39] S. Tan, Z. Wang, S. Cheng, Z. Liu, J. Han, W. Fu, Processing maps and hot workability of Super304H austenitic heat-resistant stainless steel, Mater. Sci. Eng., A 517 (2009) 312–315. [40] X. Li, Y. Zou, Z. Zhang, Z. Zou, Microstructure evolution of a novel Super304H steel aged at high temperatures, Mater. Trans. 51 (2010) 305–309. [41] Y. Zhang, L. Zhu, A. Qi, Z. Lu, Microstructural evolution and the effect on mechanical properties of S30432 heat-resistant steel during aging at 650 °C, ISIJ Int. 50 (2010) 596–600. [42] J. Jiang, L. Zhu, Strengthening mechanisms of precipitates in S30432 heatresistant steel during short-term aging, Mater. Sci. Eng., A 539 (2012) 170– 176. [43] K. Zhan, C.H. Jiang, V. Ji, Effect of prestress state on surface layer characteristic of S30432 austenitic stainless steel in shot peening process, Mater. Des. 42 (2012) 89–93. [44] Y. Gao, C.L. Zhang, X.H. Xiong, Z.J. Zheng, M. Zhu, Intergranular corrosion susceptibility of a novel Super304H stainless steel, Eng. Fail. Anal. 24 (2012) 26–32. [45] T. Amadou, C. Braham, H. Sidhom, Double loop electrochemical potentiokinetic reactivation test optimization in checking of duplex stainless steel intergranular corrosion susceptibility, Metall. Mater. Trans. A 35A (2004) 3499–3513. [46] T. Sourmail, T. Okuda, J.E. Taylor, Formation of chromium borides in quenched modified 310 austenitic stainless steel, Scr. Mater. 50 (2004) 1271–1276. [47] A.P. Majidi, M.A. Streicher, The double loop reactivation method for detecting sensitization in AISI 304 stainless steels, Corrosion 40 (1984) 584–593. [48] B. Deng, Y.M. Jiang, J.L. Xu, T. Sun, J. Gao, L.H. Zhang, W. Zhang, J. Li, Application of the modified electrochemical potentiodynamic reactivation method to detect susceptibility to intergranular corrosion of a newly developed lean duplex stainless steel LDX2101, Corros. Sci. 52 (2010) 969–977. [49] J. Gong, Y.M. Jiang, B. Deng, J.L. Xu, J.P. Hu, J. Li, Evaluation of intergranular corrosion susceptibility of UNS S31803 duplex stainless steel with an optimized double loop electrochemical potentiokinetic reactivation method, Electrochim. Acta 55 (2010) 5077–5083. [50] J. Hong, D. Han, H. Tan, J. Li, Y. Jiang, Evaluation of aged duplex stainless steel UNS S32750 susceptibility to intergranular corrosion by optimized double loop electrochemical potentiokinetic reactivation method, Corros. Sci. 68 (2013) 249–255. [51] F. Ruel, P. Volovitch, L. Peguet, A. Gaugain, K. Ogle, On the origin of the second anodic peak during the polarization of stainless steel in sulfuric acid, Corrosion 69 (2013) 536–542. [52] K.S. de Assis, F.V.V. de Sousa, M. Miranda, I.C.P. Margarit-Mattos, V. Vivier, O.R. Mattos, Assessment of electrochemical methods used on corrosion of superduplex stainless steel, Corros. Sci. 59 (2012) 71–80. [53] K.S. de Assis, A.C. Rocha, I.C.P. Margarit-Mattos, F.A.S. Serra, O.R. Mattos, Practical aspects on the use of on-site double loop electrochemical potentiodynamic reactivation technique (DL-EPR) for duplex stainless steel, Corros. Sci. 74 (2013) 250–255. [54] Y.S. Sato, H. Kokawa, Preferential precipitation site of sigma phase in duplex stainless steel weld metal, Scr. Mater. 40 (1999) 659–663. [55] E.E. Stansbury, R.A. Buchanan, Fundamentals of Electrochemical Corrosion, ASM International, 2000. [56] K.H. Lo, C.T. Kwok, W.K. Chan, Characterisation of duplex stainless steel subjected to long-term annealing in the sigma phase formation temperature range by the DLEPR test, Corros. Sci. 53 (2011) 3697–3703.