Influence of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel welding joints

Influence of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel welding joints

Accepted Manuscript Title: Influence of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel welding jo...

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Accepted Manuscript Title: Influence of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel welding joints Author: Zhiqiang Zhang Hongyang Jing Lianyong Xu Yongdian Han Lei Zhao Jianli Zhang PII: DOI: Reference:

S0169-4332(16)32163-8 http://dx.doi.org/doi:10.1016/j.apsusc.2016.10.047 APSUSC 34141

To appear in:

APSUSC

Received date: Revised date: Accepted date:

14-9-2016 6-10-2016 7-10-2016

Please cite this article as: Zhiqiang Zhang, Hongyang Jing, Lianyong Xu, Yongdian Han, Lei Zhao, Jianli Zhang, Influence of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel welding joints, Applied Surface Science http://dx.doi.org/10.1016/j.apsusc.2016.10.047 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Influence of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel welding joints

Zhiqiang Zhanga, b, Hongyang Jinga, b, Lianyong Xua, b *, Yongdian Hana, b, Lei Zhaoa, b, Jianli Zhangc

a

School of Materials Science and Engineering, Tianjin University, Tianjin 300350, China

b c

Tianjin Key Laboratory of Advanced Joining Technology, Tianjin 300350, China

Welding laboratory, Offshore Oil Engineering (Qing Dao) Company, Qing Dao 266520, China

*Corresponding author. Tel./Fax: +86 22 27402439 E-mail address: [email protected]

1

GRAPHICAL ABSTRACT:

2

Highlights 3



N2-supplemented shielding gas promoted nitrogen solid-solution in the austenite.



Secondary austenite had higher Ni but lower Cr and Mo than primary austenite.



Pitting corrosion preferentially occurred at secondary austenite and Cr2N.



Adding N2 in shielding gas improved pitting corrosion resistance of GTAW joint.



E2209T1 weld metal had very poor pitting corrosion resistance due to inclusions.

Abstract: The influences of microstructure and elemental partitioning on pitting corrosion resistance of duplex stainless steel joints welded by gas tungsten arc welding (GTAW) and flux-cored arc welding (FCAW) with different shielding gas compositions were studied by optical microscopy, electron backscatter diffraction, scanning electron microscopy, transmission electron microscopy, energy

dispersive

spectroscopy,

electron

probe

microanalysis,

and

potentiostatic

and

potentiodynamic polarization methods. The adding 2% N2 in shielding gas facilitated primary austenite formation in GTAW weld metal (WM) and suppressed Cr2N precipitation in GTAW weld root. In the HAZ, the banded microstructure disappeared while the coarse ferrite grains maintained same orientation as the banded ferrite in the BM. In the WM, the ferrite had one single orientation throughout a grain, whereas several families of austenite appeared. The austenite both in BM and WM enriched in Ni and nitrogen, while Cr and Mo were concentrated in the ferrite and thus no element showed clear dendritic distribution in the WM (ER2209 and E2209T1). In addition, the secondary austenite had higher Ni content but lower Cr and Mo content than the primary austenite. The N2-supplemented shielding gas promoted nitrogen solid-solution in the primary and secondary austenite. Furthermore, the secondary austenite had relatively lower pitting resistance equivalent number (PREN) than the ferrite and primary austenite, thereby resulting in its preferential corrosion. The Cr2N precipitation led to relatively poor resistance to pitting corrosion in three HAZs and pure Ar shielding GTAW weld root. The N2-supplemented shielding gas improved pitting corrosion resistance of GTAW joint by increasing PREN of secondary austenite and suppressing Cr2N precipitation. In addition, the FCAW WM had much poorer resistance to pitting corrosion than the GTAW WM due to many O-Ti-Si-Mn inclusions. In the BM, since the austenite with lower PREN 4

compared to the ferrite, the pitting corrosion occurred at the ferrite and austenite interface or within the austenite.

Keywords: duplex stainless steel; welding; microstructure; elemental partitioning; electron probe microanalysis; pitting corrosion

1 Introduction Duplex stainless steel (DSS) with nearly equal proportion of ferrite () and austenite () phase combines most of the beneficial properties of austenitic stainless steel (ASS) and ferritic stainless steel (FSS), including excellent mechanical properties, well corrosion resistance, good formability and weldability, and lower cost than the ASS due to the addition of low amounts of Ni [1-3]. Therefore, DSS was increasingly being used for various applications such as submarine pipelines and riser system in petroleum and natural gas industries, seawater cooled heat exchanger tubing, pressure vessels and boilers, desalination facilities, and nuclear power facilities [4-8]. To attain excellent properties, DSS is recommended to maintain a ferrite/austenite (/) ratio close to 1:1 without secondary phase precipitation [9]. However, this dual-phase balance in weld metal (WM) and heat-affected zone (HAZ) will be disturbed during welding [10,11]. It is well known that welding is an inevitable fabrication process in most of applications of DSS. Gas tungsten arc welding (GTAW) is one of the most popular techniques for welding DSS because it produces high-quality welds, although relatively low efficiency restricts its wide applications [12-14]. In recent years, flux-cored arc welding (FCAW) attracts many attentions in welding DSS due to its higher welding efficiency compared to GTAW [11]. However, the studies on microstructure and corrosion resistance of DSS FCAW joint are limited in the current literatures. In both HAZ and WM, the microstructure undergoes rapid heating and cooling cycles, thus resulting in excessive ferrite and some undesirable secondary precipitates such as Cr2N, M23C6, and sigma-phase () [14-19]. Excessive ferritization and undesirable secondary precipitates can cause catastrophic deterioration of 5

mechanical properties and corrosion resistance of the DSS joint [20]. The Cr2N is one of the most common precipitates in HAZ. The Cr2N precipitation forms a Cr-depleted zone in the surroundings, which ultimately results in the localized corrosion [21,22]. In addition, when the WM and HAZ in this metastable condition is reheated, as in multi-pass welding, the most apparent changes is secondary austenite (2) precipitation [17,23-25]. The 2 can improve the toughness of DSS but compromises its localized corrosion resistance [20]. It has been reported by Chenhua et al [26] that the localized corrosion of DSS welds is associated with the 2 due to the lower Cr and Mo contents. However, the further in-depth studies on pitting corrosion behavior associated with the 2 and Cr2N precipitation have not been proposed in current technical documents and literatures. Nitrogen is added to DSS in order to promote austenite formation and further solid solution in austenite for improving its mechanical properties and corrosion resistance [27-31]. However, nitrogen loss is inevitably during welding process [31]. The nitrogen addition in shielding gas is a relatively effective and economic method to compensate for nitrogen loss in DSS WM and HAZ for ensuring desirable microstructure and properties [22,28]. To date, some studies have reported on the effects of nitrogen alloying on the microstructure evolution and properties of DSS WM and HAZ. Westin and Johansson et al [28,32,33] have reported that nitrogen loss from the weld pool can lead to a deterioration of pitting corrosion resistance due to high ferrite contents and associated Cr2N precipitation, thus the nitrogen in the shielding gas and backing gas can increase significantly localized corrosion resistance of DSS WM. Tsuge et al [34] studied the influence of nitrogen content on localized corrosion resistance of simulated DSS welding joint, and indicated that the increase of nitrogen content obviously improved pitting corrosion resistance and stress corrosion cracking resistance because of an improvement in corrosion resistance of austenite phase by the dissolved nitrogen in itself. It was also reported by Kim et al. [22] that pitting corrosion resistance of DSS after welding with N2-supplemented Ar shielding gas was greatly improved due to the reduction of the Cr2N. However, the effects of nitrogen added in Ar shielding gas on elemental partitioning in the different phases and pitting corrosion resistance in the GTAW WM and HAZ are yet to be investigated. The aim of this work was to investigate the influences of microstructure and elemental 6

partitioning in different phases on pitting corrosion resistance of DSS joints welded by GTAW and FCAW techniques in different shielding gas compositions by using optical microscopy (OM), electron backscatter diffraction (EBSD), scanning electron microscopy (SEM), transmission electron microscopy (TEM), energy dispersive spectroscopy (EDS), electron probe microanalysis (EPMA), and potentiostatic and potentiodynamic polarization methods. 2 Experimental procedure 2.1 Materials and welding procedure A 14 mm thickness DSS plate (UNS S31803) was used as base metal (BM) in this study, which had been hot-rolled followed by solution treatment at 1040–1100 °C for 1 h and then quenched in water. The GTAW and FCAW techniques were utilized in preparing DSS joints. According to WRC (Welding Research Council) diagram [35], the ER2209 and E2209T1, both with 1.2-mm diameters, were used as filler wires for the GTAW and FCAW, respectively. Karlsson et al [36,37] also reported that 22% Cr standard DSSs are welded with the well-established 22Cr9Ni3Mo+N type matching filler materials. The chemical compositions of the BM and the different wires are listed in Table 1. As compared to the BM, both filler wires had higher Ni and Cu contents for facilitating austenite formation. To comparatively study effects of N2 addition in the Ar shielding gas, the shielding gas compositions used in GTAW contained pure Ar and 98% Ar + 2% N2 mixed gas at 20 L/min. In addition, 100% CO2 shielding gas at 15 L/min was applied for FCAW. The backing gas with nitrogen is expected to have good influence on an improvement of corrosion resistance on the rootside, as reported by Westin et al [28]. However, the pure Ar used as backing gas in this work is mainly aimed to study the influence of N2 addition in the Ar shielding gas during GTAW process on microstructure, elemental partitioning, and pitting corrosion resistance. In the future, we will do special study the influence of backing gas with nitrogen on it. The schematic diagram of the welding process is shown in Fig. 1. To simulate special joints such as tee and corner joints in DSS special application, a single bevel with an angle of 45 was used. This single bevel had a root gap of 2–4 mm and a root face of 1–2 mm. The backing gas of pure Ar at 20 L/min was applied in the process of root welding, and was maintained until the third pass was completed. A total of 18 passes were needed for the GTAW (weld root: 3 passes and weld filler: 15 passes), whereas only 13 passes were required for the FCAW (root 7

welded by the GTAW in the pure Ar shielding: 3 passes and FCAW weld filler: 10 passes). The heat input (Q, KJ/mm) can be expressed as follows:

Q 

U *I S

(1)

where I: average welding current (GTAW in pure Ar: 135.9 A, GTAW in 98% Ar + 2% N2: 138.6 A, and FCAW: 187.2 A); U: average welding voltage (GTAW in both pure Ar and 98% Ar + 2% N2: 12.6 V, and FCAW: 24.8 V); S: average welding speed (GTAW in pure Ar: 2.7 mm/s, GTAW in 98% Ar + 2% N2: 2.0 mm/s, and FCAW: 4.7 mm/s); η: welding thermal efficiency (GTAW: 0.65 and FCAW: 0.9). The average heat input during the GTAW welding in pure Ar and 98% Ar + 2% N2 shielding and FCAW in 100% CO2 shielding were 0.61, 0.62, and 0.88 KJ/mm, respectively. The interpass temperature was controlled at <150C. 2.2 Microstructure characterization Cross sectional specimens of the GTAW and FCAW joints [see Fig. 1(b)] were ground successively from 280, 400, 600, 800, 1000, and 1500 to 2000 grit and then polished using diamond pastes (particle size: 2.5, 0.5, and 0.25 µm). To distinguish different austenite phases and small precipitates such as Cr2N, the as-polished specimens were electrochemically etched in a 10 wt.% oxalic acid solution for 15–20 s at 0.4 A direct current. The microstructure was characterized by OM (ZEISS, Axio Vert.A1) and SEM (ZEISS, EVO 18) coupled with EDS (OXFORD, X-MaxN). The chemical compositions of the different phases after electrochemically etching were quantitatively measured by EPMA instrument (FE-EPMA, JXA-8530F). The quantitative calibrations of metallic elements (including Fe, Cr, Mo, Ni, Si, Mn, and Cu) and nitrogen were performed by the corresponding high pure metals and boron nitride, respectively. To reduce the random error, the least 20 times measurements for each phase in the different zones were performed and then the average values were calculated as final results. In addition, after mechanically polished to mirror finish followed by electrolytic polished with a 98 vol.% ethyl alcohol and 2 vol.% perchloric acid (HClO 4) solution applying 45 V at 20 °C for 20 s, three different weld center metals were used for mapping by the EPMA. The instrument contained five wavelength dispersive spectrometers for mapping. The grain orientation analysis was done by EBSD (OXFORD, NordlysMax2). Transverse cross section (RD-TD plane) specimens [see Fig. 1(b)] for EBSD analysis were prepared by using same procedure 8

as the EPMA mapping. The austenite contents in the different regions were calculated by using the OM with Pro-Image software. The OM coupled with Pro-Image software provided a nice functional module to calculate different phase contents. The measured regions (including weld root, weld center, HAZ, and BM) were shown schematically in Fig. 1(b). 10 OM images in each zone were collected at 500 magnification and then were processed, including grayness transform, binarization, and phase extraction and calculation, finally the average value was calculated. The small precipitates such as Cr2N were further characterized by TEM (FEI, Tecnal G2 F30) coupled with EDS. 2.3 Electrochemical testing To evaluate the resistance pitting corrosion in the different zones of the DSS joints welded by three different techniques, potentiostatic and potentiodynamic polarization measurements were performed in a three-electrode cell with a potentiostat (Gamry, Interface 1000). A platinum foil, saturated calomel electrode (SCE), and tested specimen were used as counter electrode, reference electrode, and work electrode, respectively. All potentials presented in this paper refer to SCE. The tested specimens (exposed area: 0.196 cm2) include weld root, weld center, mixed zone (WM + HAZ + BM), and BM. To reduce the effect of surface condition on test result, all specimens were successively ground from 280 to 1,500 grit using a series of emery papers and mechanically polished to mirror finish. 1 M (mol/L) NaCl was chosen as the test solution, which was purged with high pure N2 at 0.3 L/min for 30 min for deoxidization before each test. Before each electrochemical test, the specimen was cathodically polarized at -0.75 VSCE for 5 min to improve reproducibility. Critical pitting temperature (CPT) of the WM and BM was first obtained by the potentiostatic polarization according to ASTM G 150 [38]. The potentiostatic test was started by applying a constant potential of 0.75 VSCE and increasing solution temperature at a rate of 1 °C/min from 25 °C until the current densities exceed 100 A/cm2 continually where the temperature is defined as the CPT. The higher the CPT, the material had better resistance to pitting corrosion. Besides, the potentiodynamic polarization was performed according to ASTM-G5 [39]. The solution temperature was kept at 60 C (near or slightly above CPT). After the open circuit potential reached a stable state (about 30 min), the test was started at a scanning rate of 0.5 mV/s from -0.5 VSCE + corrosion potential (Ecorro) to pitting potential (Epit) at which the current density indicated that stable pitting or transpassivity had 9

occurred. The Epit was defined as the potential where the current density continuously exceeded 100 μA/cm2. The potential difference of Epit - Ecorr was used to evaluate pitting corrosion resistance. In general, the materials with larger Epit - Ecorr values exhibited better pitting corrosion resistance due to more stable passivation film formation [40]. To ensure reproducibility, each electrochemical experiment was performed in triplicate. Finally, pits morphologies were observed using the SEM. 3 Results and discussion 3.1 Microstructure characterization Fig. 2 shows the optical microstructures in the different zones of the DSS joint welded by three different welding techniques. In the DSS (S31803) welding joints, the austenite contains two types: primary austenite and 2. The primary austenite are directly solidified from molten metals among ferritic dendrites as well as formed from the subsequent solid-state phase transition of the ferrite into the austenite. According to precipitation mechanism and morphology, the primary austenite are divided into grain boundary austenite, Widmanstätten austenite, intragranular austenite, and partially transformed austenite. The allotriomorphic austenite first begun to nucleate and grew at prior ferrite boundaries, which were named as grain boundary austenite [41,42]. When cooling continued and the amount of grain boundary austenite increased, the available nucleation sites at the ferrite boundaries decreased and then if time permitted, new nuclei begun to initiate from grain boundary austenite, and grew towards the ferrite side in the form of side-plate Widmanstätten austenite [41]. If with sufficient time after reheating from subsequent beads, the intragranular austenite can form at the Ni/N-rich regions in the ferrite grain. The intragranular austenite formation needs more undercooling as a driving force compared to the grain boundary austenite and Widmanstätten austenite because of its higher activation energy for lattice diffusion, as reported by Eghlimi et al [42]. Since lattice diffusion occurs much slower than grain boundary diffusion, the intragranular austenite growth is limited, thereby resulting in finer grains, as shown in Fig 2. The intragranular austenite generally nucleates at the inclusions and dislocations or other small precipitates such as Cr2N [24,25]. Previous study shows that the ferrite gain size is a critical factor to control the intragranular austenite formation [25]. Hertzman, Brolund, and Wessman et al [43,44] also thought that a smaller grain size and cooling rate is vitally to enhance austenite formation. In addition, most of the austenite is expected to dissolve 10

after reheating to the dual-phase region with the coexistence of the ferrite and austenite. However, a small portion of residual austenite remained in the ferrite, named partially transformed austenite, which appeared as a plateau structure (see Fig. 2). The residual partially transformed austenite had important influences on effectively inhibiting the Cr and Mo segregation at the solidus temperature in addition to preventing the growth of the ferrite grain [42]. Furthermore, the DSS HAZ contains coarse ferrite grains with insufficient austenite content in addition to some secondary precipitates. Generally, the HAZ is divided into the high temperature HAZ (HTHAZ) and the low temperature HAZ (LTHAZ). According to previous report, the temperature for DSS LTHAZ is approximately in the range of 950–650°C, while the peak temperature of the HTHAZ reached above 1000 °C [32]. The microstructures in the LTHAZ are similar to the BM and do not form secondary precipitates. Thus, most studies are performed at the HTHAZ because of apparent structural changes resulting in deterioration of the toughness and corrosion resistance [20]. Figs. 3-5 show the EBSD images of the DSS joint welded by three different welding techniques. The BM presented a typical banded structure of alternating austenite and ferrite elongated in the rolling direction [45]. The Euler maps [see Figs. 3(a)-5(a)] revealed one ferrite or austenite band consisted of several different grains with different orientations. Especially, the austenite grains are fragmented indicating a partial or complete recrystallization process after solution treatment, and few annealing twins are visible in the austenite. The inverse pole figure (IPF) in the rolling direction [see Figs. 3(c)-5(c)] and in the normal direction [see Figs. 3(d)-5(d)] showed that most of the ferrite grains presented a rolling texture component {001}〈110〉, while the austenite grains showed random orientation, as reported by Badji et al [45]. In the HTHAZ, the banded microstructure disappeared and the equiaxed ferrite grain formed. However, the coarse ferrite grains in the HAZ maintained the same orientation as the banded ferrite in the BM. The allotriomorphic austenite and intragranular austenite precipitated along the ferrite boundary and within the ferrite grain in the HAZ, respectively. The fine intragranular austenite was mainly randomly oriented, whereas coarser austenite was close to Kurdjumov–Sachs (KS) orientation relationship (OR), as reported by Karlsson et al [46]. In the WM, the ferrite has one single orientation throughout a grain, whereas several families of austenite appeared. Allotriomorphic grain boundary austenite had a random OR with one of the adjacent 11

ferrite grains and was close to KS with the other, whereas Widmanstätten austenite always had an OR close to KS with the matrix. In the future, the recrystallized degree, special boundary, misorientation, and texture in the different zones of the DSS welding joint will be further studied based on the EBSD. Fig. 6 shows the austenite contents in the weld root, weld center, HAZ, and BM of the DSS joints welded by three different techniques. During the GTAW process, the N2-supplemented shielding gas increased the austenite contents in the different zones (weld root: 35.2%  54.3%, weld center: 39.2%  56.4%, and HAZ: 23.5%  31.0%). Fig. 2 reveals that 2% N2 addition in Ar shielding gas mainly promoted the primary austenite formation and increase the primary austenite grain size. The austenite content in the FCAW weld center reached 49.3%, which was significantly higher than that in the GTAW weld center with the pure Ar shielding but lower than that in the GTAW weld center with the 98% Ar + 2% N2 shielding. The austenite content in the GTAW weld root with the pure Ar shielding was increased from 35.2% to 41.6% by the subsequent FCAW process. The FCAW HAZ presented a slightly higher austenite content (27.1%) than the GTAW HAZ (23.5%) in the pure Ar shielding but lower value than the GTAW HAZ in the 98% Ar + 2% N2 shielding (31%). Two aspects can be used to explain why the WM and HAZ in the FCAW joint presented higher austenite contents compared to those by the GTAW in the pure Ar shielding: (a) higher heat input during the FCAW than the GTAW resulting in a more sufficient transformation of the ferrite to the austenite; (b) each weld bead received more filler metal during the FCAW than the GTAW, resulting in a longer dwell time at the temperature where austenite formation during reheating from the subsequent FCAW process. Therefore, the Euler maps in Figs. 3(b)-5(b) shows that the FCAW HAZ exhibited a wider range than the GTAW HAZ, and the equiaxed ferrite and austenite grains formed due to complete recrystallization process. However, three HAZs exhibited relatively low austenite content because of the high reheating temperature and the rapid cooling. The previous study reported that the peak temperature in the HAZ is close to 1350 °C where the coexistence of the ferrite phases and the small amount of liquid phase, and then the HAZ rapidly cools resulting in an insufficient austenite content [19]. Hertzman, Brolund, and Wessman et al [43,44] studied that austenite reformation in the HAZ of three duplex stainless steels 12

(molybdenum-free 2304, 2205, and super duplex 2507) by using experimental simulation and theoretical calculation methods. They indicated that the cooling rates, grain size and alloy content were three key factors to control the reformed austenite content. In the multi-pass welding, underlying and side WM as well as HAZ are reheated by deposition of the subsequent passes and then, the apparent changes of microstructure in these reheated zones are 2 and Cr2N precipitation [24,25]. There are two different types of 2 observed in the WM and HAZ by the SEM, as shown in Figs. 7-9. The first type is the intragranular 2, which nucleates and grows intragranularly in the ferrite. However, it is hard to distinguish the intragranular primary austenite and intragranular 2. Generally, the intragranular 2 have smaller size than the intragranular primary austenite. The second type is the intergranular 2, which precipitates at the interface of the ferrite and primary austenite and grows towards the ferrite side. The etching technique formed a clear difference when the intergranular 2 had grown from the primary austenite, which can be easily observed by the SEM. In the HAZ, the Cr2N were prone to precipitation at the austenite-free ferrite regions or along the ferrite grain boundary as well as at the /2 interface [see Fig. 7(e) and (f), 8(e) and (f), and 9(e) and (f)]. In addition, there were small amounts of Cr2N found in the GTAW weld root with the pure Ar shielding [see Fig. 7(b)]. Nitrogen loss is inevitably during welding process [31]. The nitrogen addition in shielding gas can compensate for nitrogen loss to promote the austenite formation [22,28]. It has been verified by the Ramirez et al [24] that the excessive ferrite supersaturated with nitrogen is crucial for Cr2N precipitation. The nitrogen loss in combination with rapid cooling rate resulted in excessive supersaturated ferrite formation in the pure Ar shielding GTAW weld root and thus led to Cr2N precipitation. The N2-supplemented shielding gas promoted the austenite formation (see Fig. 6) and therefore decreased the supersaturated ferrite content, and finally suppressed the Cr2N precipitation in the GTAW weld root. In addition, the reheating from higher heat input and more deposits of the subsequent FCAW beads led to the Cr2N dissolution in the pure Ar shielding GTAW weld root. Furthermore, the Cr2N was not observed in the weld center. There are also two types of Cr2N: the intragranular and intergranular Cr2N, as revealed by the SEM in Fig. 7(b), (e) and (f), 8(e) and (f), and 9(e) and (f) and the TEM in Fig. 10(a), (c), and (d). The Cr2N presented a rod-like morphology. The electron diffraction patterns reveal that the Cr2N belongs to close-packed 13

hexagonal structure. The EDS line analysis in the TEM [see Fig. 10(b)] shows that the enrichment of Cr in the Cr2N resulted in the Cr depletion in the ferrite adjacent to the Cr2N, as reported by Kim et al [22]. Most of the Cr2N were formed intragranularly and along the ferrite grain boundaries. However, some Cr2N were found within the 2 and at the /2 boundaries. It was reported by Ramirez et al [24,25] that the predominant orientation relationship of the //Cr2N was

110α 111 γ  0001Cr N (the most compact planes of three phases were parallel) and 2

 111α  110 γ  1100Cr N , which resulted in the formation of low-energy interfaces, thus making the 2

Cr2N precipitation at the / interfaces energetically favorable. Furthermore, a great number of inclusions were found in the FCAW WM, as shown in Figs. 2, and 9(c) and (d). The EDS analysis revealed that chemical compositions of the inclusions in the FCAW WM consisted of O, Ti, Si, and Mn, as shown in Fig. 11. 3.2 Elemental partitioning in different phases Thermodynamic calculation shows both BM (UNS S31803) and WM (ER2209 and E2209T1) belongs to ferritic-austenitic solidification mode and thus, it is not possible to form a microstructure with 100% ferrite [20]. However, the solidification occurs with ferrite as the primary phase in the HTHAZ and WM due to rapid cooling after welding [53]. Fig. 12 shows the element distribution in the BM. The ferrite stabilizers such as Cr and Mo were significantly enriched in the ferrite, while the austenite stabilizers such as Ni and nitrogen were concentrated in the austenite, as generally accepted. However, the partitioning of metallic elements in the WM especially for low-nitrogen DSS is not as apparent as in the BM, while nitrogen is generally enriched in the austenite, as reported by Westin and Ogawa [47,48]. Westin and Hertzman et al [47] have studied the element partitioning by EPMA mapping LDX 2101 DSS weld and revealed that segregation of metallic elements could follow the dendritic solidification structure. Generally, the ferritic solidification in the DSS welds occurs as epitaxial growth of the ferrite grains from the fusion boundary and the dendrite growth is controlled by the thermal gradient [32]. This dendritic solidification structure could lead to particular segregation of metallic elements. In this segregation condition, the Cr is evenly distributed in the WM, while the Ni, Mn, Mo, and Si are enriched in the interdendritic area. Nitrogen is generally 14

enriched in the austenite in the WM. By comparing EPMA mapping results of WMs (22Cr–6Ni–3Mo–0.12N, 22Cr–6Ni–3Mo–0.18N, 22Cr–10Ni–3Mo, and 22Cr–10Ni–3Mo-0.12N), Ogawa et al [48] drew a conclusion that nitrogen alloying can significantly increase the austenite fraction and promote the partitioning of the Ni and nitrogen in the austenite, while the Cr and Mo in the ferrite. The chemical compositions of the ER2209 (22Cr–9Ni–3Mo–0.15N) and E2209T1 (22Cr–9Ni–3Mo–0.146N) used in this study are same to and even slightly even higher nitrogen content than 22Cr–10Ni–3Mo-0.12N WM studied by Ogawa. Thus, it is expected that the ferrite dendritic structure is not as apparent as reported by Ogawa and Westin et al. Figs. 13-15 show the different element distributions in the GTAW WM and FCAW WM. It can be clearly seen that the austenite enriched in Ni and nitrogen, while the ferrite enriched in Cr and Mo, and at the same time, no element showed clear dendritic distribution. However, the partitioning of the Ni, Mo, and nitrogen in the ferrite and austenite was more distinct than the Cr partitioning in the WM. The metallic element partitioning in the FCAW WM was more obvious than that in the GTAW WM. In addition, both in the GTAW and FCAW WM, the Cr and Mo elements were more concentrated in the ferrite between the austenite than within the large ferrite grain. Tables 2-4 show that the Widmanstätten austenite has lower amounts of Cr, Mo and nitrogen than the grain boundary austenite, although enriched in Ni, because it forms at lower temperatures and after consumption of most of the alloying elements by grain boundary austenite, as reported by Eghlimi et al [42]. The intragranular austenite exhibited slightly lower Cr and Mo contents and higher Ni content than other primary austenite. In addition, the 2 had higher Ni content but significantly lower Cr and Mo contents than the primary austenite, as shown in Figs 13-15, and Tables 2-4. However, the same as the primary austenite, the 2 enriched in the nitrogen compared with the ferrite. Furthermore, it can be clearly seen that the adding 2% N2 in shielding gas significant promoted nitrogen solid-solution in the primary austenite and 2, while there was little influence on nitrogen solid-solution in the ferrite, as reported by Ogawa et al [48]. Previous study has reported that availability and diffusion of nitrogen is crucial for formation and growth of austenite [31]. Westin et al [32] have reported that a higher nitrogen content results in more efficient austenite reformation and makes the alloy less sensitive to rapid cooling rates. 3.3 Cooperative precipitation mechanism of 2 and Cr2N 15

Under the solution-annealed condition, the DSS retains almost all the nitrogen in solid solution within the austenite because of the high solubility and low diffusion of nitrogen in the austenite as compared to that in the ferrite [20]. Thus, as the ferritized microstructure after welding cools rapidly from the high temperature, the ferrite would become supersaturated in nitrogen. When this metastable structure in the WM and HAZ are reheated, as in the multi-pass welding, the 2 and Cr2N precipitation would compete with each other [24,25]. In general, the Cr2N is not a stable phase at the temperature where the 2 precipitation occurs [24]. However, there were still many Cr2N remained in the metastable microstructure after reheating to high temperature in the multi-pass welding due to short holding time and rapid cooling rate. Based on the detailed microstructure characterization and chemical composition analysis, a mechanism for the cooperative precipitation of the 2 and Cr2N is proposed, as shown schematically in Fig. 16. First, the Cr2N nucleates heterogeneous at the supersaturated ferrite in nitrogen, which is facilitated by the low-energy interfaces [24]. The Cr2N growth will cause the Cr depletion and Ni enrichment in the surrounding ferrite, which generates the favorable composition conditions for the 2 precipitation. Therefore, the 2 usually has lower Cr content and higher Ni content than the primary austenite, as listed in Tables 2-4. It is expected that there is a large supply of nitrogen from the supersaturated ferrite [24]. The sufficient nitrogen via the long diffusion path is support for the Cr2N growth towards the ferrite side. However, the diffusion of heavy elements such as Cr and Ni is the limiting factor for the Cr2N growth because of their short diffusion path. Then, the 2 nucleates in the Cr2N (see Fig. 17) and grows towards the supersaturated ferrite. In addition, some intergranular 2 can directly attach to the primary austenite and then grows towards the supersaturated ferrite side, which does not require nucleation. When reheating is done at temperatures above 950 °C, the Cr2N left behind [see Fig. 7(b)] become unstable and dissolve, as reported in literature [24]. Thus, the Cr from the Cr2N dissolution could enrich in the core of the 2 due to its short diffusion path while the nitrogen can diffuse into the 2 because of the long diffusion path and the high solubility. Unfortunately, this phenomenon was not evidenced by the EPMA because of detection limit. Namely, the actual Cr content in the 2 maybe lower than the EPMA results in Tables 2-4. In addition, the nitrogen from the Cr2N dissolution can participate in the 2 precipitation. The 2 growth draws the Ni from the surrounding ferrite and rejects the Cr to the ferrite, 16

which generates the favorable condition for the Cr2N precipitation again at the /2 boundary, as shown in Fig. 7 (e) and (f), 8(e) and (f), and 9 (f). 3.4 Pitting corrosion resistance Fig. 18 illustrates the typical curves of the current density against temperature obtained from the potentiostatic polarization test in the deaerated 1 M NaCl solution. It is clearly observed that the value of the current density is lower than 10 μA/cm2 during the initial heating process, which implies that the specimen surfaces are protected well by the passive film. When the electrolyte temperature increases continuously up to the critical pitting temperature, a sharp rise of current density is found, implying an occurrence of some stable pits. The average CPT obtained from the potentiostatic polarization tests are listed in Table 5. The BM exhibited the highest CPT (57.3 °C). The 2% N2 added into the Ar shielding gas during the GTAW significantly increased the CPT of the weld root metal (43.8 °C48.5 °C) and the weld center metal (47.5 °C51.9 °C). In addition, the FCAW WM (45.7 °C) presented relatively lower CPT than the GTAW weld center metal (pure Ar: 47.5 °C and 98% Ar+2% N2: 51.9 °C). The subsequent FCAW reheating process increased the CPT of the pure Ar shielding GTAW weld root metal (43.8 °C46.6 °C). Fig. 19 shows the potentiodynamic polarization curves in the weld root, weld center, WM + HAZ + BM, and BM of three different welding joints in the deaerated 1 M NaCl solution at 60 C. It is appropriate that 60 C was used as potentiodynamic polarization temperature, which should be equal to or slight above CPT according to ASTM-G5. The average values of the Ecorr, Epit, and Epit – Ecorr obtained from the potentiodynamic polarization tests are listed in Table 6. The potential difference of Epit - Ecorr was used to evaluate the pitting corrosion resistance of the different zones in the DSS joint welded by three different techniques. The mixed zone exhibited relatively low Epit Ecorr in comparison with the WM and BM, thus indicating a poor resistance to pitting corrosion in the HAZ. The N2-supplemented shielding gas significantly increased Epit - Ecorr of the different zones in the GTAW joint (weld root: 0.5631 VSCE0.9272 VSCE, weld center: 0.9134 VSCE0.9749 VSCE, and WM + HAZ + BM: 0.4773 VSCE0.7733 VSCE). The FCAW WM (0.6503 VSCE) exhibited relatively lower Epit - Ecorr than the GTAW WM. In addition, the BM had the highest Epit - Ecorr (0.9973 VSCE), which indicated an excellent resistance to pitting corrosion as compared to the DSS 17

welding joints. Figs. 20-22 demonstrate the SEM pits morphologies in the different zones of three different welding joints after the potentiodynamic polarization test in the deaerated 1 M NaCl solution at 60 C. It can be clearly seen that the pitting corrosion occurred at the 2 or the /2 boundary in the GTAW weld root and weld center, as shown in Fig. 20(a), (c) and (d), 21(a)-(d), and 22(a) and (b). In the GTAW weld root with pure Ar shielding, a large amount of stable pits were observed in the ferrite adjacent to the Cr2N [see Fig. 20(b)] in addition to at the /2 boundary [see Fig. 20(a)]. With the addition of 2% N2 in the Ar shielding gas, the Cr2N precipitation in the weld root was suppressed. Consequently, the pitting corrosion degree in the 98% Ar + 2% N2 shielding GTAW weld root was much slighter than the pure Ar shielding GTAW weld root. In addition, the FCAW WM presented relatively severe pitting corrosion around the inclusions, as shown in Fig. 22(c)-(d). In the HAZ, a grent number of pits were observedin the ferrite grain adjacent to the Cr2N, as shown in Figs. 20(e), 21(e) and (f), and 22(e) and (f). In addition, similar to the GTAW WM, the 2 in the HAZ were prone to pitting corrosion compared to the ferrite and primary austenite, as shown in Figs. 20(f)-22(f). Fig. 23 shows the pits morphologies of the BM. There were many metastable and stable pits found within the austenite and at / boundary. A common method to rank the resistance to pitting corrosion is using a pitting resistance equivalent number (PREN). In general, the higher PREN represented better resistance to pitting corrosion. The PREN is related to three of the most essential elements (Cr, Mo and nitrogen), which is calculated by each element weighted according to the contribution to the resistance to pitting corrosion, as follows [11,47]:

PREN=wt.% Cr + 3.3  wt.% Mo + 16  wt.% N

(2)

The PREN values of different phases in the different zones of three different DSS joints are shown in Tables 2-4. The 2 in the WM and HAZ had relatively lower PREN value than the ferrite and primary austenite, thereby resulting in preferential corrosion of 2, which was consistent with observation of the corroded surface. In addition, the 2 in the WM exhibited lower PREN value than that in the HAZ, which maybe because of lower Cr and Mo contents in addition to more nitrogen loss from the weld pool during the welding. The N2-supplemented shielding gas increased significantly the PREN of 2 as well as primary austenite and promoted the formation of the primary austenite with relatively high 18

PREN value, thereby led to a great improvement of pitting corrosion resistance in the GTAW WM and HAZ. In the pure Ar shielding GTAW weld root and three HAZs, although with higher PREN than the 2, the ferrite phases and the / boundary were still attacked around the Cr2N due to Cr depletion in the ferrite adjacent to the Cr2N. However, the C2N precipitation in the GTAW weld root was significantly suppressed by adding 2% N2 in the Ar shielding gas, which was one of the most important reasons for the increase of Epit - Ecorr in the GTAW weld root. In addition, although the 2 in the FCAW WM with higher PREN compared to the GTAW WM in the pure Ar shielding, the FCAW WM had much poorer resistance to pitting corrosion, which is mainly because of a large amount of O-Ti-Si-Mn inclusions resulting in the initiation and propagation of the pits. Therefore, the inclusions must be controlled to a relatively level before the FCAW method is applied to weld DSS in the corrosive condition. In the BM, since the austenite with lower PREN compared to the ferrite, the pitting corrosion occurred at the / interface or within the austenite. 4 Conclusions (1) The N2-supplemented shielding gas facilitated the primary austenite formation in the GTAW WM and suppressed the Cr2N precipitation in the GTAW weld root. The FCAW joint presented a higher austenite content than the pure Ar shielding GTAW joint but lower than the 98% Ar + 2% N2 shielding GTAW joint. And many O-Ti-Si-Mn inclusions were distributed in the FCAW WM. (2) In the HAZ, the banded microstructure disappeared and the equiaxed ferrite grain formed. However, the coarse ferrite grains maintained the same orientation relationship as the banded ferrite in the BM. In the WM, the ferrite had one single orientation throughout a grain, whereas several families of austenite appeared. (3) The austenite in both the BM and WM enriched in Ni and nitrogen, while Cr and Mo were concentrated in the ferrite, and thus no element showed clear dendritic distribution in the WM. The 2 had higher Ni content but significantly lower Cr and Mo content than the primary austenite. However, the same as the primary austenite, the nitrogen was enriched in the 2 compared to the ferrite. Furthermore, the adding 2% N2 in shielding gas significantly promoted nitrogen solid-solution in the primary austenite and 2, while there was little influence on nitrogen solid-solution in the ferrite. (4) Based on the detailed microstructure characterization and chemical composition analysis, a 19

mechanism for the cooperative precipitation of the 2 and Cr2N is proposed. Since the Cr-depleted/Ni-rich zone formation in the ferrite around the Cr2N, the Cr2N precipitation generated the favorable constitutional conditions for the 2 formation, and at the same time, acted as the nucleation sites for the 2. In addition, the 2 growth caused the Ni depletion and Cr enrichment in the ferrite, thereby resulting in the Cr2N precipitation again at the /2 boundary. (5) The HAZ exhibited relatively poor resistance to pitting corrosion in comparison with WM and BM. The N2-supplemented shielding gas significantly improve the pitting corrosion resistance of the GTAW joint. The FCAW WM had much poorer resistance to pitting corrosion than the GTAW WM. The BM showed excellent resistance to pitting corrosion compared to the DSS welding joints. (6) The 2 in the WM and HAZ had relatively lower PREN value than the ferrite and primary austenite, thereby resulting in preferential corrosion of 2. In the pure Ar shielding weld root and three HAZs, although with higher PREN than the 2, the ferrite phases and the / boundary were still attacked due to the Cr depletion in the ferrite around the Cr2N. The N2-supplemented shielding gas increased the PREN of 2 as well as primary austenite. The FCAW WM had much poorer resistance to pitting corrosion, which is mainly because of a large amount of inclusions resulting in the initiation and propagation of the pits. In the BM, since the austenite with lower PREN compared to the ferrite, the pitting corrosion occurred at the / interface or within the austenite.

Acknowledgments This work was supported by the National Natural Science Foundation of China [grant number 51575382]; and Marine Economic Innovation and Development of Regional Demonstration Projects of China [Grant number cxsf2014-12].

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Figure captions

Fig. 1 The schematic diagram of the welding process, tested zones in multi-pass welding joint, and EBSD examined surface.

Fig. 2 The optical microstructures in the different zones of the DSS joint welded by three different welding techniques. 26

Fig. 3 The EBSD images of the DSS GTAW joint in the pure Ar shielding: (a) phase map presenting the ferrite in blue and the austenite in yellow; (b) Euler map showing the orientation of different ferrite and austenite grains; (c) IPF in the rolling direction (see triangle for the color code); (d) IPF in the normal direction (see triangle for the color code).

Fig. 4 The EBSD images of the DSS GTAW joint in the 98% Ar + 2% N2 shielding: (a) phase map presenting the ferrite in blue and the austenite in yellow; (b) Euler map showing the orientation of different ferrite and austenite grains; (c) IPF in the rolling direction (see triangle for the color code); (d) IPF in the normal direction (see triangle for the color code).

27

Fig. 5 The EBSD images of the DSS FCAW joint: (a) phase map presenting the ferrite in blue and the austenite in yellow; (b) Euler map showing the orientation of different ferrite and austenite grains; (c) IPF in the rolling direction (see triangle for the color code); (d) IPF in the normal direction (see triangle for the color code).

Fig. 6. The austenite contents in the weld root, weld center, HAZ, and BM of the DSS joints welded by three different techniques.

28

Fig. 7 The 2 and Cr2N characterization by the SEM in the different zones of the DSS GTAW joint in the pure Ar shielding.

Fig. 8 The 2 and Cr2N characterization by the SEM in the different zones of the DSS GTAW joint in the 98% Ar + 2% N2 shielding.

29

Fig. 9 The 2 and Cr2N characterization by the SEM in the different zones of the DSS FCAW joint but root welded by the GTAW in the pure Ar shielding.

30

Fig. 10 The Cr2N characterization by the TEM: (a) intragranular Cr2N within the ferrite grain; (b) the line analysis of Cr content adjacent to Cr2N; (c) intergranular Cr2N at the / boundary; (d) intergranular Cr2N at the /2 boundary.

Fig. 11 The EDS analysis of inclusions in the FCAW WM.

31

Fig. 12 The distribution of different elements in the BM.

Fig. 13 Distribution of different elements in the GTAW WM with the pure Ar as shielding.

32

Fig. 14 Distribution of different elements in the GTAW WM with the 98% Ar + 2% N2 shielding.

Fig. 15 Distribution of different elements in the FCAW WM.

33

Fig. 16 The evolution of the 2 and Cr2N precipitation.

Fig. 17 The Cr2N acted as nucleation site of the 2.

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Fig. 18 The typical curves of the current density against temperature obtained from the potentiostatic polarization tests in the deaerated 1 M NaCl solution: (a) weld root; (b) weld center.

Fig. 19 The potentiodynamic polarization curves in the different zones of three different welding joints in the deaerated 1 M NaCl solution at 60 C: (a) weld root; (b) weld center; (c) WM+HAZ+BM.

35

Fig. 20 Pits morphologies of the different zones in the GTAW joint with the pure Ar gas shielding after the potentiodynamic polarization test in the deaerated 1 M NaCl solution at 60 C.

Fig. 21 Pits morphologies of the different zones in the GTAW joint with the 98% Ar+2% N2 shielding gas after the potentiodynamic polarization test in the deaerated 1 M NaCl solution at 60 C.

36

Fig. 22 Pits morphologies of the different zones in the FCAW joint after the potentiodynamic polarization test in the deaerated 1 M NaCl solution at 60 C.

Fig. 23 Pits morphologies of the BM after the potentiodynamic polarization test in the deaerated 1 M NaCl solution at 60 C.

37

Table captions Table 1 Chemical compositions of the base material and the different wires. Chemical composition (wt.%) Materials

C

Si

Mn

P

S

Ni

N

Cr

Mo

Cu

Fe

Base material (UNS S31803)

0.018

0.54

0.92

0.011

0.003

5.3

0.17

22.9

3.0

0.042

Bal.

Filler wire for GTAW (ER2209)

0.008

0.48

1.54

0.017

0.0006

8.63

0.15

22.94

3.07

0.14

Bal.

Filler wire for FCAW (E2209T1)

0.04

0.46

1.08

0.017

0.011

8.73

0.146

23.02

3.57

0.10

Bal.

1

Table 2 Chemical compositions of different phases in GTAW joint with pure Ar shielding. Element (wt.%) zone

Phase

PREN

Cr

Mo

Ni

N

Ferrite

23.768

3.048

7.967

0.043

34.51

Grain boundary austenite

22.963

2.781

8.443

0.293

36.83

Widmanstätten austenite

22.535

2.663

9.219

0.234

35.07

Intragranular austenite

22.351

2.635

9.318

0.120

32.97

Partially transformed austenite

23.481

2.979

7.977

0.198

36.48

Secondary austenite

22.026

2.477

9.818

0.079

31.46

Ferrite

24.497

3.829

4.544

0.039

37.76

Primary austenite

23.314

3.294

6.247

0.352

39.82

Secondary austenite

22.004

2.572

6.916

0.200

33.69

Ferrite

24.401

3.852

4.441

0.045

37.83

Austenite

21.633

2.511

6.765

0.257

34.03

WM

HAZ

BM Notes: Pitting resistance equivalent number (PREN) = wt.% Cr + 3.3 wt.% Mo + 16 wt.% N.

2

Table 3 Chemical compositions of different phases in GTAW joint with 98% Ar+2% N2 shielding. Element (wt.%) zone

Phase

PREN Cr

Mo

Ni

N

Ferrite

23.584

3.073

7.501

0.045

34.44

Grain boundary austenite

23.091

2.748

8.011

0.347

37.71

Widmanstätten austenite

23.069

2.713

8.317

0.309

36.97

Intragranular austenite

22.022

2.685

8.971

0.280

35.36

Partially transformed austenite

23.204

2.804

7.753

0.295

37.18

Secondary austenite

21.553

2.347

9.637

0.181

32.19

Ferrite

24.349

3.545

4.717

0.043

36.74

Primary austenite

22.891

2.912

5.776

0.472

40.05

Secondary austenite

22.239

2.520

6.882

0.325

35.76

WM

HAZ

Notes: Pitting resistance equivalent number (PREN) = wt.% Cr + 3.3 wt.% Mo + 16 wt.% N.

3

Table 4 Chemical compositions of different phases in FCAW joint with pure CO2 shielding. Element (wt.%) zone

Phase

PREN Cr

Mo

Ni

N

Ferrite

23.291

3.467

8.651

0.035

35.29

Grain boundary austenite

22.992

3.182

8.828

0.251

37.51

Widmanstätten austenite

22.725

3.027

9.145

0.232

36.43

Intragranular austenite

22.342

2.926

9.732

0.204

35.26

Partially transformed austenite

23.078

3.253

8.795

0.217

37.28

Secondary austenite

20.859

2.441

9.418

0.193

32.00

Ferrite

24.122

3.497

5.532

0.042

36.33

Primary austenite

23.425

3.155

6.253

0.287

38.43

Secondary austenite

22.347

2.634

6.725

0.231

34.74

WM

HAZ

Notes: Pitting resistance equivalent number (PREN) = wt.% Cr + 3.3 wt.% Mo + 16 wt.% N.

4

Table 5 The average CPT obtained from the potentiostatic polarization tests. Welding methods

Testing zone

CPT (°C)

Weld root

43.8

Weld center

47.5

Weld root

48.5

Weld center

51.9

Weld root

46.6

Weld center

45.7

BM

57.3

GTAW with pure Ar

GTAW with 98% Ar + 2% N2

FCAW but root welded by GTAW –

5

Table 6 The average values of the Ecorr, Epit, and Epit – Ecorr obtained from potentiodynamic polarization tests. Welding methods

GTAW with pure Ar

GTAW with 98% Ar + 2% N2

FCAW but root welded by GTAW



Testing zone

Ecorr (VSCE)

Epit (VSCE)

Epit – Ecorr (VSCE)

Weld root

-0.3001

0.2632

0.5631

Weld center

-0.3047

0.6087

0.9134

WM+HAZ+BM

-0.2995

0.1778

0.4773

Weld root

-0.3316

0.5956

0.9272

Weld center

-0.3034

0.6715

0.9749

WM+HAZ+BM

-0.2847

0.4886

0.7733

Weld root

-0.2971

0.5071

0.8042

Weld center

-0.3272

0.3231

0.6503

WM+HAZ+BM

-0.3326

0.2216

0.5542

BM

-0.2831

0.7142

6

0.9973