Corrosion Science 53 (2011) 3756–3763

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Corrosion Science

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The effect of large heat input on the microstructure and corrosion behaviourof simulated heat affected zone in 2205 duplex stainless steel

Yinhui Yang ⇑, Biao Yan ⇑, Jie Li, Jia WangSchool of Materials Science and Engineering, Tongji University, Shanghai 200092, PR ChinaShanghai Key Lab of D&A for Metal-Functional Materials, Shanghai 200092, PR China

a r t i c l e i n f o a b s t r a c t

Article history:Received 28 December 2010Accepted 16 July 2011Available online 22 July 2011

Keywords:A. Stainless steelB. Thermal cyclingC. Pitting corrosionC. Intergranular corrosion

0010-938X/$ - see front matter � 2011 Elsevier Ltd.doi:10.1016/j.corsci.2011.07.022

⇑ Corresponding authors at: School of Materials SciUniversity, Shanghai 200092, PR China. Tel.: +86 21853 85.

E-mail address: yanbiao_tj@163.com (B. Yan).

In the simulated heat affected zone of 2205 duplex stainless steels, effects of large welding heat inputs onthe microstructure and corrosion behaviour were investigated. Reformed austenite content increasedwith the coarsening of grain boundary austenite (GBA) and the growth of intragranular austenite (IGA)and Widmanstatten austenite (WA), thus improving the low temperature toughness and affecting corro-sion state. Reduction of chromium nitrides contributed to better resistance to pitting corrosion. More-over, the pitting corrosion and intergranular corrosion were improved resulting from the formation ofmore GBA and WA. The specimen with a Dt8/5 of 100 s presents better comprehensive performance.

� 2011 Elsevier Ltd. All rights reserved.

1. Introduction

The application of a steel as structure material is highly influ-enced by its weldability. Thus, it is of great significance to improvewelding efficiency without sacrificing the performance of weldedjoint. The implementation of large welding heat input to large-sizeworkpieces has always been an effective way to promote produc-tivity, e.g. the connection of oil pipelines. However, among thethree parts of a welded joint – the base metal, the weld metaland the heat affected zone (HAZ), the strength and the corrosionresistance of the HAZ are the most easily degraded properties be-cause of the intermetallic precipitation and grain coarsening. Atpresent, the low carbon steels are commonly used as pipelinematerials, whose coarse grain HAZ (CGHAZ) can lead to degradedmechanical properties [1].

By comparison, duplex stainless steels (DSSs) which containboth ferrite and austenite phases exhibit better mechanical proper-ties and corrosion resistance, as well as better weldability andlower sensitivity to the weld cracking. This enables various typesof DSSs to be used in certain fields such as oil and gas, chemical,paper, petrochemical, desalination, and power generation [2–6].

2205 DSS (UNS S 31803) is one of the most common kinds ofDSSs with the volume fraction of each phase above 30%. The goodmechanical property and corrosion resistance depend largely uponthe proper austenite–ferrite balance [7,8]. Nevertheless, during

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welding, the phase ratio (c/d) in either the fusion zone or theHAZ tends to deviate from 1:1, which considerably influencesmechanical properties and corrosion resistance. After the homoge-nized microstructure was obtained at the welding peak tempera-ture, the rapid microstructural changes can be induced by thephase transformation during the cooling. Particularly, it is reportedthat the microstructure of the HAZ in DSSs presents degradedchemicals and impact properties compared with that of the finelysolubilized steels [9,10]. Meanwhile, the slow cooling rate (withlarge heat input) can lead to favourable high reformed austenitecontent in DSSs, thus affecting the corrosion behaviour and somemechanical properties of the simulated HAZ [11].

Therefore, it is crucial to improve the performance of HAZ withappropriate heat input. Although certain amount of research workhas been carried out on the microstructure and pitting corrosion inthe HAZ of DSSs [12,13], no systematic build-up on the effects oflarge heat input on a microstructure and corrosion behaviour inthe HAZ of DSSs has been proposed. It is not clear whether DSSsare suitable for high current and large heat input welding as well.Therefore, the aim of this paper is to investigate the effect of largeheat input on the microstructure and corrosion behaviour of thesimulated HAZ in 2205 DSSs, and obtain an appropriate heat inputrange for improving the welding efficiency.

2. Experimental method

2.1. Materials and heat treatment

The material investigated was a commercial standard duplexstainless steel 2205 plate (UNS31803) with the composition shown

Table 1Chemical composition (wt.%) of the alloy used.

C Mn Si S P Ni Cr Mo N Fe

0.023 1.35 0.38 0.001 0.026 4.83 22.07 2.37 0.19 Bal.

Y. Yang et al. / Corrosion Science 53 (2011) 3756–3763 3757

in Table 1. Prior to weld simulation, the steel plate was subjectedto a solution treatment at 1060 �C for 30 min, and then quenchedin water to get the homogenized structures. The specimens, withdimensions of 11 � 11 � 80 mm, were cut from a 12 mm-thick-plate with the grains elongated in the rolling direction. The Gleeble3800 thermo-mechanical simulator was used for the welding HAZsimulation, and its thermal cycles are schematically shown inFig. 1. It can be observed that the cooling rate will be reduced withlonger cooling time. The specimens were heated to a peak temper-ature of 1345 �C at the rate of 400 �C/s, held for 1 s and then cooledat various rates to simulate different heat inputs. It is considered atypical temperature range of 1200–800 �C for the formation of aus-tenite and precipitation in DSSs. However, the cooling time from800 to 500 �C (Dt8/5) is relatively easier to be measured accuratelythan that from 1200 to 800 �C (Dt12/8). Thus, we can calculate theDt12/8 by measuring the Dt8/5 merely and following the Eq. (1) [13],

Dt8=5

Dt12=8¼

1ð500�T0Þ2

� 1ð800�T0Þ2

1ð800�T0Þ2

� 1ð1200�T0Þ2

ð1Þ

where T0 = starting temperature (�C).Therefore, the Dt8/5 was preferably tested to represent the cool-

ing time of welding thermal cycle. The relationship between theDt8/5 and heat input follows the Eq. (2) [6].

Qd¼ kðDt8=5Þ1=2 ð2Þ

where Q is heat input (J mm�1), d is the thickness (11 mm), k is thethermal coefficient (25.52 J mm�2 s�1/2) for DSSs. As the Dt8/5 in-cludes 20, 50, 80, 100 and 120 s, the corresponding heat inputsare 1.26, 1.98, 2.51, 2.81 and 3.08 kJ mm�1, respectively.

2.2. Microstructural characterization and corrosion resistance tests

The thermal cycle samples, used for microstructural observa-tion and corrosion resistance tests, were cut along the section ofthe thermocouple position. To observe the microstructure, themechanically polished samples longitudinal to the rolling direction

Fig. 1. Thermal cycles of Gleeble simulation.

were electrolytically etched in 40 wt.% KOH solutions at 6 V for15 s, and then examined by optical microscope (OM). To identifythe precipitates, the scanning electron microscopy (SEM PhillipsXL30 FEG) with an energy dispersive X-ray (EDX) detector and atransmission electron microscope was used. The TEM thin foilswere prepared by jet polishing in a mixture of 10% perchloric acidand 90% methyl alcohol at 253 K and 15 V. Specimens of10 � 10 � 55 mm were manufactured for Charpy impact test (V-notch) at �40 �C. Phase volume fractions were obtained using themethod of manual point count according to ASTM E 562 as follows[14]: the magnification of the micrograph was 500� and the gridsize (number of points) was 25. Any point that fell on the phasestudied was counted as one, otherwise zero. In addition, the pointfell on the boundary was counted as a half. The ferrite grain sizewas measured by the linear intercept method. Then, the corrosionresistance evaluation was carried out by means of anodic potentio-dynamic polarization measurement in a deaerated 3.5 wt.% NaClsolution at 30 �C, which was comprised of three electrodes. A plat-inum foil and a saturated calomel electrode (SCE) were used as thecounter and reference electrodes, respectively. The specimens,embedded in epoxy resin with an exposure area of 100 mm2, actedas a working electrode. Prior to each experiment, the specimenswere ground mechanically up to 2000 grit, rinsed with distilledwater and dried in hot air. Anodic potentiodynamic polarizationwas tested through linear sweep technique at a sweep rate of0.1 mV/s, from the free corrosion potential to 1200 mV(SCE). Allpotentials are given against the SCE. After that, SEM was used toobserve corrosion morphology.

The susceptibility to intergranular corrosion (IGC) was evalu-ated by Double-loop electrochemical potentiokinetic reactivation(DL-EPR) test, which was conducted in a solution containing1 M H2SO4, 0.5 M NaCl and 0.01 M KSCN at 30 ± 1 �C at a scan rateof 1 mV/s. The KSCN and NaCl were used to break the passive filmduring the reactivation cycle of the tests [15], and the three-elec-trode cell was also used in the DL-EPR test. The potential isscanned in the anodic direction from corrosion potential (Ecorr)to a point of 0.250 V(SCE) in the middle of the passive region.Next, the scanning direction is reversed and the potential is re-duced back to the cathodic region, so an anodic loop and a reac-tivation loop are generated. The maximum current density (Ia) inthe anodic scan loop and the maximum current density (Ir) in thereversed scan loop were measured simultaneously. The degreeof sensitization (Ra) was measured by the given formula of(Ir/Ia) � 100%.

3. Results and discussion

3.1. Microstructural characterization

Fig. 2 shows the optical microstructure of the specimen afterthe solution treatment at 1060 �C for 30 min. The austenite phase(c) is shown as white while the gray region is the d-ferrite phase(d). As shown in this figure, the island-like c phase was surroundedby the continuous matrix of the d-ferrite phase. The volume frac-tion counting results showed that the specimens contained 53.2%c and 46.8% d-ferrite after the solution treatment, indicating a goodbalance of both phases.

Fig. 3 shows the effect of the cooling time (Dt8/5) on the micro-structural changes in the simulated HAZ of 2205 DSSs. At 1345 �C,most of c phases were dissolved while d-ferrite phases kept stableafter being held for 1 s. Besides, a little amount of the c phase re-mained, referred to as partially transformed austenite (PTA), hadsome effects on hindering the growth of d-ferrite grains [16]. Witha Dt8/5 of 20 s, the grain boundary austenite (GBA) was initiallyformed in the interfaces between d-ferrite grains, and a little

Fig. 2. Optical microstructure of 2205 DSS solution treated at 1060 �C for 30 min.

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intragranular austenite (IGA) appeared within the d-ferrite grains(Fig. 3a). In addition to GBA and IGA, the Widmanstatten austenite

Fig. 3. Optical microstructure of the HAZ of 2205 DSS after different cooling time(e) Dt8/5 = 120 s.

(WA) along the d-ferrite grain boundaries began to grow into theinterior, and the IGA became coarser when the Dt8/5 was increasedto 50 s (Fig. 3b). Both the amount of WA and IGA increased and theGBA coarsened much more at a much slower cooling rate (Dt8/5 =80 s, 100 s), as shown in Fig. 3c and d respectively. As the Dt8/5

was increased to 120 s, it can be observed explicitly that a largeamount of WA and IGA formed within d-ferrite grains, and the vol-ume fraction of reformed austenite reached 44.5% by manual pointcount measurement. Thus, the amount of reformed austenite (GBA,WA and IGA) increased by increasing welding heat input. This canbe understood by that prolonging the cooling time would increasethe transformation time, so the diffusion of c-stabilizing elementssuch as nitrogen and nickel was enhanced and more austenite wastransformed from d-ferrite.

Fig. 4 shows the TEM micrographs and diffraction patterns asthe Dt8/5 ranged from 20 to 120 s. It was reported that Cr2N isthe main precipitate that was formed in the simulated HAZ of DSSs[17,18]. However, in this study, two types of precipitates were ob-served in simulated HAZ with a Dt8/5 of 20 s, which were identifiedas tetragonal-like Cr2N (Fig. 4a and b) and platelet-like CrN (Fig. 4cand d), respectively by diffraction pattern analysis. It was also

treatments: (a) Dt8/5 = 20 s; (b) Dt8/5 = 50 s; (c) Dt8/5 = 80 s; (d) Dt8/5 = 100 s;

Fig. 4. The transmission electron microscopy graphs of the HAZ: (a) The tetragonal-like Cr2N precipitate and (b) its diffraction pattern; (c) The tiny platelet-like CrNprecipitate and (d) its diffraction pattern; (e) CrN and Cr2N precipitates with a Dt8/5 of 20 s; (f) The morphology of ferrite matrix with a Dt8/5 of 120 s.

Y. Yang et al. / Corrosion Science 53 (2011) 3756–3763 3759

observed that the amount of CrN precipitates was higher than thatof Cr2N, and these nitrides was formed primarily on the ferritephase, with little amount on the austenite phase. Fig. 4e showeda large amount of CrN and Cr2N precipitates appeared on the ferritephase with a Dt8/5 of 20 s. As the Dt8/5 was increased to 120 s, it isapparently shown in Fig. 4f that such nitrides are relatively scarce,which may be related to the reformation of more austenite at amuch slower cooling rate.

3.2. Effect of cooling time on the ferrite grain size and austenitecontent in simulated HAZ

The d-ferrite grain size as a function of the Dt8/5 is schematicallypresented in Fig. 5a. It can be seen that the d-ferrite grain size in-creased by prolonging the Dt8/5, and the larger heat input was ap-plied to the specimens, the coarser d-ferrite grains became. It isnoted that longer dwelling time of peak temperature and slowercooling rate both contribute to larger d-ferrite grain size in the

simulated HAZ. However, in this experiment, where the dwellingtime was short as 1s, thus the cooling rate became the dominantfactor to the grain coarsening. With a Dt8/5 of 80 s, the ferrite grainsize was 40 lm larger than that with a Dt8/5 of 20 s, and as theDt8/5 was increased to 120 s, the ferrite grain size was large as298 lm. According to the Hall–Petch relationship, the grain sizeis inherently relevant to the strength of steel [19]. Therefore, thestrength of HAZ is degraded to some extent due to coarsening ofd-ferrite grains.

The effect of different cooling time on the austenite content insimulated HAZ is shown in Fig. 5b. With a Dt8/5 of 20 s, the austen-ite volume fraction was 30.9% and was consisted mainly of GBA.When the Dt8/5 was increased to 80 s, the WA and IGA grewquickly, which led to a rapid increase in the amount of reformedaustenite. As the Dt8/5 prolonged to 100 s, the growth of austenitecontent got slower due to slower growth of WA and IGA. When theDt8/5 was increased to 120 s, the austenite content became higherbecause of the coarsening of WA and the formation of more IGA

Fig. 5. The effect of cooling time on the (a) ferrite grain size and (b) austenite content in simulated HAZ.

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within d-ferrite grains, which was attributed to longer cooling timecorresponding to larger heat input [15].

3.3. Effect of cooling time on the impact toughness and hardnessvariation

The microhardness value of the ferrite and austenite phases insimulated HAZ as a function of the Dt8/5 is plotted in Fig. 6a. Themicrohardness value ranged from 230 to 260 HV. With a Dt8/5 of20 s, the microhardness value of the ferrite was high in large partdue to a high d-ferrite content. Moreover, a relative large amountof nitrides precipitates within ferrite grains also accounted forthe high microhardness value of the ferrite [3,20]. By prolongingthe Dt8/5, GBA coarsened and the amount of WA and IGA increasedrapidly. This lead to a decrease in the microhardness value. In com-parison, the microhardness value of the austenite phase remainedstable around 225 HV, slightly lower than that of the ferrite phase[21].

The impact energy (Akv) as a function of Dt8/5 is shown inFig. 6b. As can be seen, the impact toughness was improved withlonger Dt8/5 on the whole, which is indicated by ever-increasingAkv. In detail, when the Dt8/5 was between 20 and 100 s, the Akv

kept a good linear relationship with the Dt8/5. Moreover, the Akv

went up markedly after a Dt8/5 of 100 s, which highly matchesthe curvilinear trend in the Fig. 5b. This could be interpreted bythe fact that increasing the amount of reformed austenite leadsto the high value of Akv with a prolonging Dt8/5.

Fig. 6. The effect of different Dt8/5 on the (a) microhard

3.4. X-ray diffractions and electron dispersive X-ray spectroscopyanalyses

Fig. 7 presents the spectra for as-received and other sampleswith different cooling time. As shown in Fig. 7a, the major peakintensities of d-ferrite and c did not differ significantly, indicatinga good balance of the two phases in the as-received specimens. Thec ? d transformation finished quite well at the welding peak tem-perature of 1345 �C, leaving nearly complete ferritic microstruc-ture by the rapid heating [22]. Then, with a Dt8/5 of 20 s (Fig. 7b),the d ? c transformation proceeded but the major peak intensityof the c phase decreased rapidly compared to the d-ferrite phaseat a high cooling rate. Moreover, the major peak intensity of thec phase was enhanced due to a prolonged cooling time of 50 s(Fig. 7c). At the lowest cooling rate, where the Dt8/5 was 120 s(Fig. 7d), the major peak intensities of d/c phases varied slightlywith larger heat inputs, indicating a relatively large volume frac-tion of reformed austenite was obtained.

Table 2 lists the chemical compositions of GBA, WA, IGA and d-ferrite determined by EDS in the simulated HAZ. A common meth-od to rank the pitting susceptibility in stainless steels is using thepitting resistance equivalent number (PREN). The PREN is linked tothe content of three of the most essential elements Cr, Mo and N,each of them weighted according to its influence on pitting:PREN = % Cr + 3.3% Mo + 30% N. This new formula was employedin this investigation instead of the traditional one (% Cr + 3.3%Mo + 16% N) for the following reasons. Nitrogen is almost com-pletely solutionized in the austenite phase whereas it is rarely

ness and (b) impact energy (Akv) in simulated HAZ.

Fig. 7. X-ray diffraction patterns in the HAZ of 2205 DSS (a) as-received specimen; and the specimens treated for different cooling time: (b) Dt8/5 = 20 s; (c) Dt8/5 = 50 s; (d)Dt8/5 = 120 s.

Table 2The chemical compositions (wt.%) and PREN values of different phases with Dt8/5 inthe HAZ.

Dt8/5 Phase Cr Ni Mo Na PREN

20 GBA 21.86 5.71 2.26 0.19 35.02WA 21.78 5.76 2.12 34.48IGA 20.01 5.82 1.95 32.15Ferrite 21.85 4.57 2.71 0.05 32.29

50 GBA 21.79 5.72 2.25 0.19 34.92WA 21.68 5.68 2.09 34.28IGA 20.02 5.85 2.02 32.39Ferrite 21.92 4.56 2.79 0.05 32.63

80 GBA 21.81 5.74 2.23 0.19 34.87WA 21.7 5.63 2.11 34.36IGA 19.99 5.88 2.03 32.39Ferrite 21.88 4.48 2.7 0.05 32.29

120 GBA 21.79 5.69 2.28 0.19 35.01WA 21.66 5.61 2.32 35.01IGA 19.98 6.01 2.03 32.38Ferrite 21.85 4.41 2.68 0.05 32.19

a Nitrogen in ferrite is taken as the saturation value of 0.05%.

Y. Yang et al. / Corrosion Science 53 (2011) 3756–3763 3761

solutionized in the d-ferrite phase in DSSs. The solubility of nitro-gen in the ferrite phase of DSSs has a maximum value of0.05 wt.%. The addition of nitrogen can change the partitioningcoefficients for chromium and molybdenum between d-ferriteand c phases. As a result, nitrogen increases the PREN of c-phasein DSSs considerably and thus increases the resistance to pittingcorrosion of the c phase, compared with that of the d-ferrite phase.For these reasons, Bernhardsson [23] proposed that the PREN coef-ficient for the nitrogen content in the DSS was almost doubled,from 16 to 30. Because rapid cooling rate enabled nitride over-saturated in ferrite, it is reasonable to reckon the content of

nitrogen in ferrite uniformly as the maximum value of 0.05 wt.%since it is hard to be measured [24]. It is noted that nearly all thec phase can transform to d-ferrite at the welding peak tempera-ture. Under this condition, the content of nitrogen was roughlyequal to the average nitrogen content in the as-received specimen.Because nitrogen is easily solutionized in austenite, the actual con-tent of nitrogen in the reformed austenite should be a little higherthan the average level. Hence, the content of nitrogen in austenitecan be reckoned as 0.19 wt.%, as the same content in the as-received specimen. As illustrated in Table 2, the content of chro-mium in different phases varied a little, with the highest in GBA,close to that in ferrite, and the lowest in IGA, which may largelyrely upon the extent of the element diffusion with different heatinput. The formation of GBA was fast so that chromium did nothave time to partition completely between d-ferrite and c phases,leaving chromium rich in GBA. On the contrary, it took rather long-er time for IGA to be formed, thus sparing more time for chromiumatoms to partition well to ferrite. By prolonging the Dt8/5, the PRENof IGA was drawing closer to that of ferrite, while the gap betweenthe PREN of GBA and WA and that of ferrite remained significant.Hence, the microstructural evolution with larger heat inputs indi-cated an improvement in pitting corrosion resistance due to thevariation of Cr content and the PREN value [25].

3.5. Potentiodynamic polarization analyses

As shown in Fig. 8, the influence of the Dt8/5 on the pitting cor-rosion resistance was evaluated by potentiodynamic polarizationmeasurements in 3.5 wt.% NaCl solution at 30 �C. The anodic polar-ization curves provide useful information concerning the potentialrange over which a material is susceptible to pitting corrosion. The

Fig. 8. Potentiodynamic polarization curves obtained for simulated HAZ in 3.5 wt.%NaCl solution at 30 �C with different cooling time treatment.

Fig. 9. The SEM morphology of pits formed on 2205 DSS after potentiodynamictest: (a) Dt8/5 = 20 s; (b) Dt8/5 = 80 s; (c) Dt8/5 = 100 s; (d) Dt8/5 = 120 s; (e) the EDSspectrum of inclusion shown in (b).

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variation in the corrosion potential (Ecorr) of different polarizationcurves was trivial, with their values around �250 mV (SCE). Itcan be observed that some small oscillation current peaks occurredbetween 80 to 280 mV (SCE), which is associated with metastableevents [26,27]. The current sharply increased with the scanningpotential, indicating the onset of pitting and the transition pointwas used to define the pitting potential (Epit). As the Dt8/5 was in-creased from 20 to 120 s, the corresponding pitting potentials (Epit)were 0.302, 0.378, 0.441, 0.446 and 0.637 V (SCE), respectively.This suggests that the Epit was shifted to nobler value with largerheat inputs and the pitting corrosion resistance was thus improvedto a certain extent. This can be explained by that larger heat inputfavored the formation of reformed austenite, thus reducing theamount of chromium nitrides precipitates like CrN and Cr2N. Inthe simulated HAZ, these nitrides were prone to precipitate withhigh ferrite content at low heat input [14], causing chromium-depleted region around them and inducing preferential sites forpitting corrosion [20,21]. On the other hand, larger heat inputsled to the growth of GBA and WA with higher PREN value, whichincreased the pitting corrosion resistance as well.

After potentiodynamic polarization tests under different Dt8/5,the SEM micrographs of pitting corrosion in the simulated HAZof DSSs are depicted in Fig. 9. As can be observed, pitting corrosionoccurred primarily within the ferrite grains with a Dt8/5 of 20 s(Fig. 9a). It is well established that ferrite is the dominating phasethrough faster cooling, which results in more nitride precipitationowing to lower nitrogen solubility than austenite [18]. Moreover,the above TEM results also showed that a large amount of CrNand Cr2N precipitation was formed within ferrite with lower heatinput. Thus, pitting corrosion could occur easily on the chro-mium-depleted region induced by the nitrides precipitation. Byprolonging the Dt8/5 to 80 s (Fig. 9b), the pits were observed onthe ferrite and IGA, also with a very little amount of compoundinclusion formed. Identified by the EDS result, the compound inclu-sion was composed of O–Mn–Al–Si–Ca (Fig. 9e), which can be ob-served both in the as-received and other HAZ specimens withdifferent heat inputs. The incoherence between the inclusionsand the adjacent region could trigger the localized corrosion ontheir interface. With a Dt8/5 of 100 s (Fig. 9c), as more reformedaustenite was formed and grew, many of pits were observed onthe IGA. When the Dt8/5 was increased to 120 s (Fig. 9), some pitswere formed on the IGA, the others were on ferrite. As indicated inTable 2, the lower PREN of IGA than GBA and WA accounted for itsdegraded pitting corrosion resistance, which was consistent with

the results from SEM micrographs analyses. In addition, a very lit-tle amount of compound inclusion was observed in the as-receivedand other HAZ specimens with different heat inputs.

3.6. Evaluation of intergranular corrosion by DL-ERP tests

DL-EPR tests were employed to evaluate the degree of sensitiza-tion of the simulated HAZ for the specimens. The lower value of thedegree of sensitization is, the better resistance to intergranular cor-rosion (IGC) is obtained. The DL-EPR curves for the specimens withvarious Dt8/5 are compared in Fig. 10. In addition, Table 3 lists thesummary of Ia, Ir, and Ra determined by the DL-EPR test results.Among these, Ia as anodic activation peak current density, was al-most the same for all the specimens, indicating that the presence ofreformed austenite did not show obvious effects on the anodic ac-tive dissolution. And the Ir values which were small at about 50 mV(SCE), decreased with longer cooling time. Therefore, the DL-EPRtest results indicated that the degree of sensitization (Ir/Ia) of thesimulated HAZ decreased by increasing heat inputs, suggestingan improved resistance to IGC. This is probably because muchmore GBA was formed at the d/d boundary interfaces by increasingDt8/5, whereas the presence of these boundary interfaces, if notcovered by GBA, would easily lead to the occurrence of IGC corro-sion. In addition, the previous EDS results indicated that theamount of GBA with high chromium content increased with longerDt8/5, presenting better recovery ability of passive layer at grain

Fig. 10. DL-EPR curves of the simulated HAZ with different Dt8/5.

Table 3The summary of Ia, Ir and Ir/Ia determined by the DL-EPR test in the simulated HAZwith different Dt8/5.

Dt8/5 (s) Ia, 10�3 A cm�2 Ir, 10�5 A cm�2 Ra = Ir/Ia (%)

20 8.342 12.62 1.5150 7.823 9.495 1.2180 8.091 8.048 0.99

100 7.685 5.146 0.67120 8.291 4.523 0.55

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boundaries. Therefore, the resistance to IGC in the simulated HAZof DSSs was improved when a larger heat input was applied tothe specimens.

4. Conclusion

(1) The GBA was initially formed with a Dt8/5 of 20 s, and coars-ened with a longer cooling time. By increasing the Dt8/5 from50 to 80 s, the amount of reformed austenite increased pro-gressively due to the growth of WA and IGA. Finally, as theDt8/5 was increased from 100 to 120 s, the amount ofreformed austenite continued to increase due to the coars-ening of the GBA and WA and the formation of more IGA.Slower cooling rate favored the reformation of more austen-ite by prolonging the dwelling time for d ? c transformationand promoting the diffusion of c-stabilizing elements suchas nitrogen and nickel.

(2) By prolonging the Dt8/5 from 20 to 120 s, the increase in d-ferrite grain size impaired the strength to some extent,meanwhile, the low temperature toughness was improveddue to more reformed austenite formation.

(3) EDS analysis showed that both GBA and WA had higherchromium contents and PREN values than IGA, indicating abetter pitting resistance.

(4) Potentiodynamic polarization tests showed that the value ofEpit became nobler due to larger heat input, indicating thatpitting corrosion resistance in the HAZ was improved. Witha Dt8/5 of 20 s, pits were formed primarily within d-ferrite.Then as large amounts of IGA were formed with longerDt8/5, more pits appeared on IGA. The formation of Cr2Nand CrN precipitates could induce chromium-depletedregions around them, thus increasing the susceptibility topitting corrosion. The amount of preferential sites for pitting

corrosion decreased due to the formation of more GBA andWA and the reduction of nitrides precipitates, which con-tributed to a better pitting corrosion resistance. DL-ERP testsindicated that the resistance to IGC in simulated HAZ of DSSswas improved by larger heat input.

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