Metallurgical and corrosion characterization of electron beam welded duplex stainless steel joints

Journal of Manufacturing Processes 50 (2020) 581–595

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Metallurgical and corrosion characterization of electron beam welded duplex stainless steel joints

T

Jastej Singh, A.S. Shahi* Department of Mechanical Engineering, Sant Longowal Institute of Engineering & Technology (Deemed to be University), Longowal, Sangrur, 148106, Punjab, India

A R T I C LE I N FO

A B S T R A C T

Keywords: Duplex stainless steel Electron beam welding Microstructure Microhardness Corrosion

Electron beam welding (EBW) process was employed autogenously to fabricate 12 mm and 18 mm thick fullpenetration butt welds of 2205 duplex stainless steel (DSS). The welds were evaluated for their metallurgical characteristics (microstructure, ferrite content and microhardness) and corrosion (intergranular and pitting) behavior in the as-welded, aged (850 °C/30 min) and aging followed by solution treated (1050 °C/30 min) conditions. The EB welds showed a relatively higher propensity to pitting corrosion than the base metal while they remained immune to the intergranular corrosion based on the degree of sensitization. Thermal aginginduced Cr and Mo-rich intermetallic σ-phase and Cr-based carbides further led to degradation of pitting resistance while their sensitization behavior remained largely unaltered. The solution treatment promoted the formation of austenite and partial dissolution of secondary phases in the welds which favored the pitting resistance restoration tendencies. This study highlights that two factors viz. microstructural phase imbalance and weld zone size would play a significant role in influencing the corrosion properties of DSS EB welds. Thus, better pitting corrosion performance could be expected from the welds fabricated using welding procedures that would result in small sized fusion zones and a lesser degree of phase imbalance.

1. Introduction Duplex stainless steels (DSSs) are two-phase (α/γ) alloys that possess combined characteristics of austenitic as well as ferritic stainless steels. The presence of austenite (γ) and ferrite (α) in comparable proportions in their microstructure imparts them an excellent combination of strength, toughness and corrosion resistance. The welding of DSSs, however, can result in an unbalanced microstructure in the fusion zone besides causing the formation of secondary phases. Such types of metallurgical transformations can have a significant influence on the properties of DSS welds. The microstructural evolution in DSSs is governed by their Creq/Nieq ratio (which is greater than 1.95) and they exhibit solidification mechanism in the single-phase ferritic mode [1]. The evolution of austenite takes from the ferrite phase via solid-state transformations. The extent of α→γ transformation in the welds is governed jointly by the elemental composition and the associated cooling rates in the weld thermal cycle. Electron beam welding (EBW) belongs to the category of low energy/high power density welding processes. It is a fusion welding process wherein heat required to melt the materials to be joined is obtained through the conversion of the kinetic energy of a concentrated ⁎

beam consisting of high-speed electrons striking the joint. One of the key advantages of this process is its capability to weld thick sections autogenously using a single weld pass. Apart from this, low levels of distortion and residual stresses, protection of the molten weld pool from atmospheric contamination by vacuum environment, process automation, etc. are the merits that distinguish this process from conventional fusion welding processes. The associated high cooling rates in the thermal cycle of EB welding, which are otherwise beneficial for the properties of austenitic stainless steels can pose certain troublesome issues in the case of duplex grades. The rapid cooling in this process can inhibit the diffusion-controlled α→γ transformation, which can result in excessively high ferrite content in the welds [2,3] and the other one is the vacuum related effusion of nitrogen from the molten weld pool [3]. The loss of N further contributes in decreasing the austenite content as it is astrong austenite stabilizing element in duplex stainless steels. To control the ferrite content/promote austenite formation in autogenous DSS EB welds, various attempts have been made by researchers like the use of remelting procedures [3] and multi-beam techniques [4] to increase the heat input (thus slow down the cooling rates and favor austenite formation). However, an increased nitrogen loss was reported with an increase in welding heat input [3]. In another

Corresponding author. E-mail addresses: [email protected], [email protected] (J. Singh), [email protected] (A.S. Shahi).

https://doi.org/10.1016/j.jmapro.2020.01.009 Received 28 February 2019; Received in revised form 15 October 2019; Accepted 6 January 2020 1526-6125/ © 2020 The Society of Manufacturing Engineers. Published by Elsevier Ltd. All rights reserved.

Journal of Manufacturing Processes 50 (2020) 581–595

J. Singh and A.S. Shahi

Fig. 1. Schematic illustration of the EB welding procedure to achieve full-penetration butt welds by incorporating an additional backing plate used in the present work.

reported that maximum austenite formation occurred at 1050 °C [14]. With the prospect of assessing the corrosion resistance degradation and the restoration capability of the joints, these temperatures were chosen for the post-weld thermal treatments in the present work. In view of limited works reported on the corrosion performance of EB welded DSS joints, a research attempt seems to be justified where corrosion behavior of thick section welds is exhaustively studied. The present work was focused on evaluating the intergranular and pitting corrosion susceptibility of electron beam welded 12 mm and 18 mm thick industrial welds of 2205 duplex stainless steel using electrochemical techniques viz. double loop electrochemical potentiokinetic reactivation (DLEPR) and Potentiodynamic anodic polarization respectively supported by investigations about their metallurgical characteristics. To facilitate exhaustive characterization of corrosion behavior of these thick section welds, a zone-wise evaluation was carried that involved selective exposure of the surface of the welds to test solution to study their local as well as gross behavior towards the corrosive environment. Moreover, the detrimental role played by aginginduced secondary precipitation in influencing their corrosion and metallurgical properties and their restoration capability in view of degraded properties after subsequent solution treatment was also investigated.

work, the use of electron beam rotations and post-heating weld procedures have been reported to influence the ferrite content in the DSS EB welds [5]. However, such procedures are yet to prove practical applicability in industrial fabrications. For the DSS welds made autogenously, excessively high ferrite content besides precipitation of Cr2N has been reported due to faster cooling rates involved in electron beam [6] and laser beam welding processes [7]. The welds were found to exhibit inferior pitting corrosion resistance than the base metal. On the other hand, Ku et al. [3] reported that EB weld with a high ferrite proportion showed superior potentiodynamic behavior than the base metal in the 1 M H2SO4 solution. Similarly, it was reported by Schmigalla et al. [8] that the high ferrite content in DSS welds had no negative effect on their pitting behavior in FeCl3 (6 wt.%) + HCl (1 wt.%) solution. The reports on the corrosion behavior of EB welded duplex steels are rather limited when compared with the conventional DSS welds. Only one work [8] could be found in the literature that reported the sensitization behavior of EB welded DSS in an electrolyte consisting of 0.5 M sulfuric acid and 0.001 M thioacetamide at 60 °C. 850 °C is the critical temperature in context with DSS 2205, exposure to which can promote the formation of various deleterious secondary phases like σ, χ, carbides and nitrides in the microstructure [9]. The composition of such phases is based on vital alloying elements such as Cr, Mo, and N, etc. which impart corrosion resisting characteristics. Such precipitation causes localized depletion of these elements around grain/interfacial boundaries, intragranular sites, thus lowering the corrosion resistance of duplex steels. It has been reported that the most significant reduction in corrosion resistance occurred at an aging temperature of 850 °C for DSS weldments [10]. An exposure time of 4 min at 850 °C has been reported to be influential enough to degrade the pitting resistance in the case of DSS [11]. The use of post-weld heat treatment (PWHT) procedures is a recommended practice and finds usefulness particularly in chemical and fossil energy industries with a requirement of high corrosion-resistant materials as per European and American standards (NORSOK MDS D42 and ASTM A928/A928M) [12]. Young et al. [13] reported the restoration of microstructural phase balance lost during laser beam welding of DSS after heat treatment at 1050 °C for 15 min. In another work, which was focused on studying the role of annealing temperatures between 1020°-1110 °C on the properties of EB welded DSS,

2. Experimental details The base material used in the present work was duplex stainless steel (DSS) grade 2205 (UNS No. S32205) in the form of 12 mm and 18 mm thick hot rolled plates. Its chemical composition in weight percentage (analyzed using an optical emission spectrometer) was: 22.338 % Cr, 3.12 % Mo, 5.429 % Ni, 0.026 % C, 0.524 % Si, 0.058 % V, 1.585 % Mn, 0.028 % P, 0.006 % Ti, 0.031 % Nb, 0.2 % N, 0.002 % W, 0.010 % S and balance as iron. Two different sets of plates were subjected to standard cleaning and polishing procedures before carrying out the welding to avoid weld defects that could arise from contamination like dust, rust, scale, etc. on the surfaces. A close square butt joint was used for both the welds. The complete penetration of the electron beam throughout the thickness in the welding procedure was achieved by using a backing plate made of DSS 2205 which was tack welded at the root side (Fig. 1) before putting the plates in the vacuum chamber. Schematic illustration of the welding 582

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Table 1 EBW process parameters used to fabricate full penetration DSS welds in the present work. Weld joint

Accelerating voltage (kV)

Beam current (mA)

Welding speed (mm/min.)

Heat input per unit length of the weld (kJ/mm)

WL (12 mm) WH (18 mm)

150 150

90 90

800 600

0.962 1.2825

Other parametric details: Focus coil current: 2205 mA, Gun to work distance: 447 mm, Root gap: < 0.25 mm, Beam oscillation pattern: Line, Oscillation frequency: 800 Hz, Beam oscillation amplitude: 2 mm and 1.8 mm for the welds WL and WH respectively, Beam offset: Nil, Gun vacuum: 1.6 × 10−6 mbar and 2 × 10−6 mbar for the welds WL and WH respectively, Chamber vacuum: 1.8 × 10−6 mbar and 5 × 10−6 mbar for the welds WL and WH respectively.

middle and bottom regions as shown in Fig. 2. The intergranular and pitting corrosion behavior of the welds was evaluated using electrochemical techniques viz. DLEPR and Potentiodynamic anodic polarization respectively. A potentiostat (Make: Gamry Instruments, Model: Reference 600) was used for electrochemical testing at room temperature. The testing system was comprised of three electrodes viz. saturated calomel electrode (SCE) as the standard electrode, graphite as the counter electrode and the test specimen as the working electrode. The use of DLEPR technique involved the following steps: (a) an initial delay of 300 s at open circuit potential before commencing the test scan (b) forward anodic activation scan starting from -0.1 V vs. open circuit value till apex potential of 500 mV and (c) cathodic reactivation scan by reversing the direction from apex potential back to the open circuit value. A scan rate of 1.67 mV/s was used for both the scans. The ratio of the maximum value of reactivation current density (Ir) in the cathodic scan to the maximum value of activation current density (Ia) in the anodic scan multiplied by 100 gave the value of the degree of sensitization (DOS) in percentage. Higher the value of DOS, the higher the susceptibility to IGC and viceversa. The specimens exhibiting DOS values > 1 % were considered to be sensitized. The standard solution of 0.5 M H2SO4+ 0.01 M KSCN meant for austenitic steels was modified and a more aggressive solution comprising 2 M H2SO4+0.01 M KSCN+0.5 M NaCl was used to detect intergranular corrosion (IGC) susceptibility in the present work [15,16]. Potentiodynamic anodic polarization measurements of the welds were carried out in a 3.5 % NaCl solution. The procedure involved an initial delay of 300 s at open circuit potential before commencing the test. The scan was initiated from an initial potential of -0.1 vs. open circuit value to an apex value of 1.5 V. Real-time potential vs. current density plots were generated by the software package DC-105 and the values of corrosion potential (Ecorr) and pitting potential (Epitt) were noted. The calculations of pit nucleation resistance (Epitt-Ecorr) were made as this parameter has been reported [17] to indicate the pitting performance of a material subjected to a corrosive environment. Higher the value, better the resistance to pitting and vice-versa. Repeatability of the electrochemical data was ensured by performing each test scan twice. Fig. 2 shows the cross-sectional views of both the welded joints, the surfaces of which were exposed to the test solution. Electrochemical testing was performed in such a manner that an exhaustive examination of the welds could be carried out in terms of covering all the regions. A total of 4 readings were taken on the 12 mm thick weld and 5 readings were taken on the 18 mm thick weld due to its larger thickness. Regions were marked with numeral notations for both the welds and the same are indicated clearly in Fig. 2. Region 1 for both the welds was comprised of completely the weld metal/ fusion zone and the rest of the regions were composite zones with varying area proportion of the weld metal and the base metal. It was further intended to examine the relative nobility of different regions of the weldment when exposed to the test solution. However, it is worth to mention that the welded structures intended for service in corrosive environments are exposed as a whole and thus gross properties are of paramount interest while dealing with the performance-related aspects of such components. Thus, special focus was made on the regions marked as 4 and 5 for the welds WL (12

procedure adopted in the present work is shown in Fig. 1. The welds were fabricated in such a manner that the backing plate formed an integral part of the joints as the electron beam was made to impinge inside the backing plate. A similar welding procedure was followed to weld both 12 mm and 18 mm thick plates and the post-weld machining facilitated the removal of backing plates. The complete details of the EB welding parameters are presented in Table 1. The calculation of the welding heat input (HI) per unit length (in kJ/mm) was done using the formula: HI= [ŋ (V × I) ×60] / [S × 1000] where ŋ is the heat source efficiency and was taken equal to 0.95, I is the beam current in mA, V is the accelerating voltage in kV and S is the welding speed in mm/min. The heat input required to fabricate 12 mm thick weld was lesser as compared to the 18 mm thick weld joint owing to the difference in the fusion requirements in the two cases. Thus, the welds were designated as WL and WH respectively and these notations have been used frequently in the subsequent discussion. The specimens for the metallurgical and corrosion studies machined out from the weldments were subjected to thermal treatments, the details of which are listed in Table 2. The first step in the procedure to carry out the isothermal treatments of the welds was the setting of microcontroller-based heating furnace to a set temperature. After attainment of the set temperature, specimens were placed in the furnace directly for the set exposure time followed by water quenching. The notations T0, T1, and T2 were used to designate the unaged, aged and solution treated (subsequent to aging) conditions respectively. The metallurgical characterization of the welds involved the microstructural examination, quantitative ferrite content measurements and microhardness evaluation besides XRD analysis. Specimens were subjected to standard grinding and polishing procedures with grit size up to 3000 followed by chemical etching in a solution comprising of 100 ml HCl+ 100 ml distilled H2O+15gm NH₄HF₂+ 1.5gm K2S2O5. This etchant was used for microstructural examination by optical microscopy (OM). For carrying out microstructural studies at higher magnification and for elemental compositional analysis, SEM equipped with EDS was used following similar metallographic procedures, but etching was done in a solution comprising of 5gm CuCl2 + 100 ml ethanol + 100 ml HCl for a reaction time of 5 min. The identification of various phases was done using XRD examination (Machine make and model: Bruker D8 Advance) with the parameters: current = 40 mA, potential = 40 kV, time per step = 0.15 s, 2θ step = 0.02°, scan range = 20°-100°. Ferrite content measurements in the welds were carried out using a Feritscope (Make: Fischer; Model: FMP30). Vickers microhardness testing was carried out using a load of 0.5 kg with a dwell time of 20 s. Microhardness and ferrite content measurements were done by covering different regions of the welds; one was moving from the top to the bottom (Fig. 2), referred to as the region along the weld centerline which involved the examination exclusively of the weld region. Another was moving across the weld centerline such that the base metal, heat affected zone (HAZ) and the weld/fusion zone were covered at the top, Table 2 Thermal treatments used in the present work. T0 T1 T2

as-welded 850 °C/30 min followed by water quench 1050 °C/30 min followed by water quench subsequent to T1 treatment

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Fig. 2. Cross-sectional views of the welded joints after removal of their respective backing plates (a) WL, 12 mm thick (b) WH, 18 mm thick. Locations labelled along the weld centerline and across the weld centerline (top, middle and bottom traverses) are the regions where the ferrite content and microhardness measurements were made. Encircled regions on the welds along with numeral notations indicate the regions which were selectively subjected to electrochemical corrosion testing involving DLEPR and Potentiodynamic anodic polarization measurements. Total 4 regions were tested in the case of WL weld and 5 regions were tested for the WH weld due to its larger thickness. Pitting corrosion testing results in terms of pit nucleation resistance (Epitt-Ecorr) are shown for different regions corresponding to the as-welded→(T0), aged (850 °C/30 min)→(T1) and aging followed by solution treated (1050 °C/30 min)→(T2) conditions respectively taken in the same order.

solution treatment at 1050 °C/30 min restored the typical morphology (Fig. 3c) similar to the unaged condition. The DSS welds follow the single-ferritic solidification mechanism wherein the melt contains the fully ferritic structure and with the initiation of solidification, ferrite to austenite transformation in the solid state takes place. Generally, the diffusion-controlled transformation of α→γ can yield austenite of different morphologies viz. allotriomorphs located at prior α/α grain boundaries, also known as grain boundary austenite (GBA), as Widmanstatten side plates (WA) growing inside the ferrite grains from GBA and lastly as intragranular austenite (IGA) precipitates. Final austenite proportion and its form are governed by the elemental composition of the molten weld pool and the cooling rates associated with the weld thermal cycle. The optical micrographs of different regions of the welds viz. WL and WH in different conditions are shown in Figs. 4 (a–l) and 5 (a–l) respectively. The top region or the reinforcement of the weld WL (Fig. 4a) was characterized by coarse ferrite grains with austenite interspersed at grain boundaries as GBA, as WA growing inside the ferrite grains from GBA and as IGA inside the ferrite grains. The region below the top (Fig. 4b) showed similar morphologies of austenite but relatively smaller ferrite grains were observed. The bottom zone showed further finer ferrite grains with austenite located at grain boundaries and intragranular sites in minute amounts with the noticeable absence of Widmanstatten side plates. A similar trend of microstructural evolution was observed for both the welds in the as-welded/unaged condition. Such types of variations observed throughout the thickness of the welds could be due to

mm thick) and WH (18 mm thick) respectively in the subsequent discussion indicative of their overall corrosion behavior. 3. Results and discussion 3.1. Macrostructure of the welds The cross-sectional views of the EB welds are shown in Fig. 2(a–b). Full-depth penetration with adequate side wall fusion was achieved with uniform weld zone formation for both the joints visibly free from any defects. Typical morphologies associated with EB welds with a wider upper and narrower bottom region were observed. The weld volume varied while moving from the top of the welds towards the root region which was the outcome of the shape of the electron beam as the heat source operating in keyhole mode welding. Since the plate thicknesses were different for both the joints, the amount of welding heat input required in both the cases was different. This variation caused significant differences in the macrostructures of both the welds in terms of dimensions of their respective fusion zones as shown in Fig. 2(a–b). 3.2. Metallurgical characterization 3.2.1. Microstructure The microstructure of the base metal in the unaged condition (Fig. 3a) showed typical bi-phase morphology with austenite bands dispersed in the ferrite matrix. Thermal aging at 850 °C/30 min resulted in distortion in the typical banded structure (Fig. 3b) and subsequent 584

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Fig. 3. Microstructure of the base metal in different conditions (a) T0, as-received (b) T1, aged at 850 °C/30 min (c) T2, solution treated at 1050 °C/30 min subsequent to aging.

(Figs. 4(e–h) and 5 (e–h)) and the regions inside the ferrite grains were affected. This observation suggests the selective dissolution of the ferrite phase due to isothermal aging. It is well known that the intermetallic precipitation is more pronounced in the ferrite phase in duplex steels owing to the faster elemental diffusion rates (approximately 100 times) in it than the austenite. Subsequent solution treatment of the welds at 1050 °C/30 min significantly influenced their microstructures (Figs. 4(i–l) and 5 (i–l)) across all the regions. The primary microstructural change that occurs during the solution treatment is that the equilibrium between austenite and ferrite is approached and its retention due to rapid cooling by quenching. The result of the same was that the austenite formation was promoted along with its coarsening for both the welds. Even the bottom regions (Figs. 4k and 5 k) which initially after welding exhibited negligible amounts of IGA, showed abundant austenite formation. The composite zones of both the welds also showed similar microstructural changes with increased austenite proportion. SEM micrographs of the aged welds along with the EDS results are shown in Fig. 6. The grain boundary austenite was identifiable along with partially transformed austenite (PTA) present inside the ferrite grains. EDS spectra 1 and 3 in Fig. 6a divulged localized enrichment of C and N suggestive of the formation of carbides or nitrides. Further, enrichment of C and N was revealed by spectrum 1 in Fig. 6b. The elemental composition of the partially transformed austenite (Fig. 6c. spectrum 1) showed enrichment of Ni with a simultaneous decrement in the percentage of Cr. The indication of the possible formation of intermetallic σ-phase was evident from spectrum 2 in Fig. 6c that divulged localized enrichment of Cr and Mo. The composition of the sigma phase is based on the Cr and Mo concentrations. It forms by eutectoid decomposition of the ferrite phase following the reaction α→ σ+γ2. The controlling step for this transformation has been reported to be the nucleation rate [18]. The temperature and the exposure duration influence the kinetics of this reaction. The nucleation and growth of the sigma phase while being enriched in Cr and Mo causes the regions of the surrounding ferrite deficit in these elements which as a result becomes unstable and transforms into austenite. The austenite which

differential cooling rates occurring at different locations owing to the shape of the electron beam operating in keyhole mode welding. In such a case, the volume of the molten metal at any instant and thus the heat content to be dissipated was more at the top region and it gradually decreased while moving towards the bottom region of the welds; thus, relatively slower cooling rates prevailed at the top region of the welds as compared to the bottom region. It is well known that GBA and WA form at relatively higher temperatures with little undercooling while IGA forms at lower temperatures and requires a higher degree of undercooling for its formation. The formation of GBA at all the locations in the welds was attributed to the enrichment of grain boundaries in austenite stabilizing elements [2] which encouraged such morphologies of austenite independent of the cooling rates prevailing at any location. While the regions at the bottom of the welds (Figs. 4c and 5 c) were characterized by a noticeable absence of WA due to relatively shorter times for its growth from GBA into the ferrite grains. Further, the IGA was also found to be in negligible amounts as compared to the upper regions of the welds. Figs. 4d and 5 d show the composite zones for both the welds including their weld zone, FBZ besides base metal. Owing to faster cooling rates involved in EB welding, no significant HAZ grain growth was observed and as such, no significant morphological changes were observed adjacent to the fusion boundary for both the welds. However, there were certain mutual differences between the microstructures of both the welds. Firstly, the proportion of austenite in the case of weld WL was observed to be more than the weld WH. Secondly, owing to lower welding heat input, and hence associated faster cooling rates in the case of weld WL, relatively finer grain structure was observed as compared to its counterpart. A relatively larger austenite formation in the weld WL could be ascribed to a larger grain boundary area per unit volume facilitated by its finer grain structure. A larger grain boundary area provided nucleation sites for grain boundary allotriomorphs and further growth of Widmanstatten side plates in the case of low heat input weld. After aging at 850 °C/30 min, the effect of secondary precipitation was observed in the microstructures of the welds. It could be seen that the austenite network at grain boundaries was almost unaffected 585

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Fig. 4. Microstructures of the top, middle, bottom and composite zones of the weld WL corresponding to different conditions (a–d) as-welded, (e–h) aged at 850 °C/ 30 min and (i–l) solution annealed at 1050 °C/30 min subsequent to aging (GBA: grain boundary austenite, WA: Widmanstatten austenite, IGA: intragranular austenite, FBZ: fusion boundary zone, BM: base metal, WM: weld metal).

along with indications of the presence of nitride of the type Cr2N. The presence of Cr2N has been reported in the welds of DSS made using processes like laser beam [7] and electron beam welding [6,14] owing to rapid cooling of the molten weld pool. Nitrogen has an excellent solubility in the ferrite phase at higher temperatures, but it experiences a drop with decrease in the temperature. Rapid cooling from higher temperatures causes precipitation of nitrides attributable to the supersaturation of N in the ferrite phase. Further, the XRD spectrum of the aged weld (Fig. 7) showed the presence of α and γ peaks, σ-phase along with carbides of the type Cr23C6 and Cr7C3 whereas the probable presence of Cr2N was also indicated. After subsequent solution treatment of the welds, dissolution of the carbides and the sigma phase was observed.

forms has an inferior composition as compared to the primary austenite nucleated from the original ferrite phase by the solid-state transformation. The secondary austenite (γ2) co-precipitates with sigma. The EDS analysis in the HAZ region of the aged weld [Spectrum 1, Fig. 6d] further suggested the possibility of nitrides and carbides formation. XRD spectra for different specimens are presented in Fig. 7. Since both the welds were made autogenously i.e. fabricated without the addition of any filler material, significant compositional differences (except in the size and morphology of the weld microstructure) in the weld metal were not expected. XRD examination was conducted for identifying various phases in the welds i.e. only qualitative assessment was to be done, so out of the two welds, the weld WL was examined, and the results so obtained could be treated to be similar for the weld WH. The XRD spectrum of the unaged base metal showed the presence of austenite and ferrite peaks with the absence of any deleterious phases. After aging, the presence of carbides of the type Cr23C6 and Cr7C3 along with the presence of intermetallic σ-phase was indicated. Furthermore, the EB weld metal showed ferrite and austenite peaks

3.2.2. Ferrite content Quantitative measurements of the ferrite content were made along and across the weld centerline for the welds and the results are presented graphically in Fig. 8(a–e). Average measurements were also 586

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Fig. 5. Microstructures of the top, middle, bottom and composite zones of the weld WH corresponding to different conditions (a–d) as-welded, (e–h) aged at 850 °C/ 30 min and (i–l) solution annealed at 1050 °C/30 min subsequent to aging (GBA: grain boundary austenite, WA: Widmanstatten austenite, IGA: intragranular austenite, FBZ: fusion boundary zone, BM: base metal, WM: weld metal).

microstructure of the weld WL was relatively finer, thus facilitating large grain boundary area per unit volume for the formation of GBA to take place. Another possible reason for lower austenite content observed in the weld WH was the use of relatively high heat input and thus a larger loss of nitrogen could be expected. This deduction is supported by the work of Krasnorutskyi et al. [3] in which the authors attempted remelting of EB welds with an intent to inject larger heat input to retard cooling rates to promote α→γ transformation, but nitrogen loss was found to increase with an increased number of remelting passes. After aging, the ferrite content of the base metal as well as the welds (WL and WH) was found to decrease and was quantitatively measured to be 37 % (range 38 ± 3 %), 45 % (range 44 ± 4 %) and 51 % (range 48 ± 7 %) respectively. This could be attributed to the decomposition of the ferrite phase into secondary precipitation owing to its metastable nature when DSSs are exposed to critical temperatures like 850 °C. After the welds were subjected to subsequent solution treatment, the ferrite content was found to increase again and was found to be 50 % (range

made to get an overall idea about the variations in the ferrite phase proportion due to welding and thermal treatments. The base metal in the as-received condition possessed an average α/γ ratio of 54/46. For the welds WL and WH, average ferrite content increased significantly and was found to be 68 % (range 67 ± 7 %) and 76 % (range 74 ± 7 %) respectively. Ferrite to austenite transformation in duplex stainless steels is diffusion-controlled in nature and slower cooling rates in general, favor such a transformation. Faster solidification of the weld pool in the EBW process inhibited the α→γ transformation which caused retention of a larger proportion of the ferrite in both the welds. Another reason in this context was the possible vacuum related loss of nitrogen from the molten weld pool [3,4]. Nitrogen is a strong austenite promoter and its reduction contributes in reducing the austenite content or the formation of excess ferrite in the DSS weld microstructure. On making mutual comparisons, the ferrite content of the weld WL was lower than the weld WH although relatively high welding heat input was used for the latter, which led to the prevalence of relatively slower cooling rates in it. This could be explained by the fact that the 587

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Fig. 6. Representative scanning electron micrographs and energy dispersive X-ray spectroscopy (EDS) results for the specimens subjected to aging at 850 °C/30 min (a, b) WHT1 and (c, d) WLT1 respectively.

across the weld centerline for the welds (Fig. 8(c–e)) divulged that the average ferrite content of the weld zones was maximum. The regions near the fusion boundary i.e. heat-affected zones (HAZs) showed lower values than their respective fusion zones. However, the HAZs possessed

50 ± 4 %), 54 % (range 53 ± 4 %) and 58 % (range 58 ± 3 %) for the base metal and the welds (WL and WH) respectively. This increase was due to the dissolution of the precipitation products like intermetallic σphase, Cr23C6 and Cr7C3 back into the matrix. The measurements made 588

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Fig. 7. XRD spectra for the DSS 2205 base metal and the weld WL for different experimental conditions.

general, the outcome of rapid cooling could also be a contributing factor. Mirakhorli et al. [19] reported that the DSS welds were harder than the base metal owing to their high ferrite proportion. Despite the variation in the ferrite content, both the welds did not differ much in their fusion zone microhardness. The weld WH had a larger proportion of ferrite phase while the weld WL was characterized by a relatively finer grain structure. After aging, the average microhardness of the welds WL and WH was found to decrease slightly and was observed to be 276 HV (range 276 ± 16) and 283 HV (range 286 ± 13) respectively. Keeping in view the variations observed throughout the welds, the aging did not influence the microhardness with significant magnitudes and a slight decrement was observed. This might be due to the low volume fraction of secondary precipitation (nitrides, carbides, and intermetallic σ-phase) induced in the welds corresponding to an aging time of 30 min. It has been observed by Calliari et al. [20] in their work involving DSS that the microhardness did not vary significantly for the low amount of

a relatively higher ferrite percentage than the base metal under a similar set of experimental conditions. 3.2.3. Microhardness The microhardness variations across different zones of the welds corresponding to different conditions are depicted graphically in Fig. 9(a–e). The base metal in the as-received condition possessed microhardness of 280 HV (range 287 ± 23). Corresponding to conditions T1 and T2, average values were measured to be 276 HV (range 281 ± 23) and 303 HV (range 311 ± 39) respectively for the base metal. The welds WL and WH in the as-welded condition possessed an average hardness (while measurements were made from the top to bottom region) of 294 HV (range 300 ± 20) and 298 HV (range 300 ± 20) respectively. Both the weld metals were harder as compared to the parent metal. This may be due to the excess ferrite phase in the microstructure of both the welds and the probable formation of nitrides. The finer grain structure encountered in the case of EB welds, in

Fig. 8. Ferrite content variations (in %) across different regions of the welds (a,b) along the weld centerline and (c,d,e) across the weld centerline for top, middle and bottom traverses respectively. 589

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Fig. 9. Microhardness variations across different regions of the welds (a,b) along the weld centerline and (c,d,e) across the weld centerline for top, middle and bottom traverses respectively.

steels imparts them complex metallurgical and precipitation characteristics. Due to that, the number of phase formations is large as compared to the austenitic steels. The phases like σ, γ, nitrides apart from carbides precipitate in the microstructure of DSSs during their exposures to high temperatures which become the cause for simultaneous enrichment and depletion of vital alloying elements along the intergranular regions causing the problem of sensitization to prevail. DLEPR curves for different specimens are shown in Fig. 10(a–f) and the quantitative data obtained from the same are presented in Table 3. The base metal in the as-received condition did not show any sensitization (Fig.10a) as indicated by the absence of any reactivation peak in the reverse cathodic scan attributable to the absence of any secondary phase and balanced microstructure. Regions 1, 2, 3 of the WL were also found to be unsensitized but when the weld as a whole (comprising WM +HAZ + BM, marked as region 4) was exposed to the test solution, very small value of DOS (0.05233 %) was registered. The weld metal (marked as region 1) of the specimen WH was also found to be unsensitized due to insignificant reactivation current density obtained in the reverse scan (Fig. 10d). The rest of the regions were also found to be unsensitized except for region 3 which registered a DOS value of 0.83 %, thus suggesting that the sensitization was not of significant magnitude. Thus, both the weld joints in the as-welded condition showed insensitivity to intergranular corrosion based on the degree of sensitization. The overall DLEPR results suggest that the welds were not severely sensitized despite the probable occurrence of nitride formation. These observations are consistent with the work of Schmigalla et al. [8] in which the EB welds of DSS were found to be unsensitized although authors reported the precipitation of nitrides and carbonitrides in the welds. Angelini et al. [18] while investigating the sensitization phenomenon on the aged 2205 DSS in 0.5 M H2SO4+ 0.5 M KSCN test solution reported that the Cr2N precipitates were not harmful to sensitization behavior due to their uniform precipitation. The Cr depletion caused by the sigma phase was stated to be detrimental for passive film stability. Moura et al. [16] in their investigation on corrosion behavior of DSS S31803 in 2 M H2SO4+0.01 M KSCN+ 0.5 M NaCl concluded that the precipitation of Cr2N did not cause sensitization but it was

secondary phases induced by aging treatments corresponding to relatively shorter durations. After subsequent solution treatment, the average microhardness values were found to increase like in the base metal and were observed to be 306 HV (range 317 ± 43) and 300 HV (309 ± 39) for the welds WL and WH respectively. This hardening could be attributed to the restoration of the microstructural phase balance and partial dissolution of secondary phases induced by the aging treatment. It was also observed that the microhardness was slightly higher towards the bottom region of the welds (Fig. 9a–b) as compared to the top region which is attributed to relatively finer grain structure encountered at these locations as discussed. While moving across the weld centerline, the microhardness of the heat affected zones (HAZs) did not reveal significant differences as compared to their respective weld metals. This was due to faster cooling rates encountered in the case of EB welding which restricted the ferrite grain growth adjacent to the fusion boundary and thus a very narrow HAZ was obtained for both the welds. 3.3. Electrochemical corrosion testing 3.3.1. Double loop electrochemical potentiokinetic reactivation (DLEPR) testing DSSs owing to their microstructural phase balance exhibit relatively superior resistance to sensitization phenomenon as compared to the conventional austenitic stainless steels. DLEPR technique is more prevalent for assessing the IGC susceptibility based on Cr-depletion at grain boundaries owing to the precipitation of carbides triggered during exposures to temperatures lying in the sensitization range. The ferrite/ austenite boundary in the case of DSSs demarcates Cr-rich (ferritic phase) and C-rich (austenitic phase) regions. Most of the Cr required for carbides formation is contributed by the ferrite phase. This leads to the formation of a wide region depleted in Cr in the ferrite while a narrower but strong Cr-depleted zone forms in the austenite owing to its relatively less compact FCC structure leading to lower elemental diffusion rates. This has been reported to contribute to intergranular corrosion (IGC) which could take place on the austenite phase boundary [21]. Further, the presence of a large number of alloying elements in duplex 590

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Fig. 10. DLEPR curves for the specimens tested in 2 M H2SO4+ 0.5 M NaCl+ 0.01KSCN solution at room temperature (a) BM and weld WL, T0 condition (b) BM and weld WL, T1 condition (c) BM and weld WL, T2 condition (d) BM and weld WH, T0 condition (e) BM and weld WH, T1 condition (f) BM and weld WH, T2 condition.

Table 3 DLEPR and Potentiodynamic anodic polarization testing results for the specimens corresponding to different conditions.

Base Metal

Weld WL (T0)

Weld WL (T1)

Weld WL (T2)

Specimen

DOS %= (Ir/ Ia)×100

Corrosion potential Ecorr (mV)

Pitting potential Epitt (mV)

Epitt- Ecorr (mV)

BT0 BT1 BT2 WLT0-1 WLT0-2 WLT0-3 WLT0-4 WLT1-1 WLT1-2 WLT1-3 WLT1-4 WLT2-1 WLT2-2 WLT2-3 WLT2-4

– – – – – – 0.05233 – – 1.3948 – – – – –

−204 −198 113 −119 −117 −202 −136 −170 −174 −177 −219 −227 −288 −313.5 −168

950.3 515.2 1044 685.8 828.7 812.2 837 480.7 549.7 223.8 350.8 708 570.4 556.6 985.7

1154.3 713.2 931 804.8 945.7 1014.2 973 650.7 723.7 400.8 569.8 935 858.4 870.1 1153.7

591

Weld WH (T0)

Weld WH (T1)

Weld WH (T2)

Specimen

DOS %= (Ir/ Ia)×100

Corrosion potential Ecorr (mV)

Pitting potential Epitt (mV)

Epitt- Ecorr (mV)

WHT0-1 WHT0-2 WHT0-3 WHT0-4 WHT0-5 WHT1-1 WHT1-2 WHT1-3 WHT1-4 WHT1-5 WHT2-1 WHT2-2 WHT2-3 WHT2-4 WHT2-5

– – 0.83727 – – – – – – 0.01138 – – – – –

−194 −238 −195 −232 −201 −204 −184 −235 −225 −258 −82.70 −427 −231 −105 −316

417.1 491 458.6 428.9 546.8 334.3 497.2 377.1 218.2 228.6 – 259.7 522.1 1025 946.8

611.1 729 653.6 660.9 747.8 538.3 681.2 612.1 443.2 486.6 – 686.7 753.1 1130 1262.8

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is more likely to be prevalent in the ferrite phase due to the solubility drop of nitrogen with decreasing temperature during weld solidification. Both the Cr and N are vital elements for resistance to pitting corrosion and their depletion in the vicinity of such precipitates favors the pit nucleation due to deterioration of the passive film while exposed to chloride environment (3.5 % NaCl in the present case). This observation is consistent with the literature in which nitride formation in electron beam [6] and laser beam welds [7] of DSS has been reported to be detrimental to their pitting resistance. On making mutual comparisons, the weld WL exhibited overall better pitting corrosion resistance which could be ascribed to its relatively balanced microstructure. On the other hand, the high heat input weld WH was characterized by a relatively smaller proportion of austenite and larger phase imbalance which could be responsible for its lower pitting resistance. Ferrite has been reported to be the phase susceptible to pitting corrosion attack in the as-welded condition [6]. It is generally accepted in the case of duplex stainless steels that the presence of microstructural phases in comparable proportions is beneficial for their overall properties including the corrosion resistance. Even though both the welds were fabricated without using any filler material, a noticeable difference in their weld metals’ pitting behavior was observed. This could be attributed to the variation in the magnitude of welding heat input used to fabricate them which eventually governs the cooling rates and hence the weld zone microstructure. This observation suggests that welding heat input in EBW can have a significant influence on the pitting resistance of the weld metal despite it being the weaker zone as compared to the base metal. Thus, the pitting resistance of the welds could be altered if there is a significant difference in the welding heat input used to fabricate them even though in autogenous mode. When the exposed surface was changed for the weld WL in such a manner that the area ratio of weld metal/base metal decreased (regions 2 and 3 respectively), the pit nucleation resistance was shifted to nobler values. When the joint as a whole (comprised of the fusion zone, heat affected zone, and parent metal) (region 4) was subjected to the test solution, pitting resistance was observed to be lower than the parent metal but higher than the zone comprising exclusively the weld metal. For the weld WH, region 1 showed the lowest pitting resistance. The region 2 which was also comprised of the base metal showed an increase in the pitting resistance as compared to the weld metal region. However, region 3 which had relatively lower weld metal/base metal area ratio did not exhibit higher pit nucleation resistance than region 2 of the weld WH. Further, region 4 which contained an increased proportion of base metal in the exposed surface showed higher pitting resistance than region 3. For the weld WH, again the pitting resistance value of the composite zones including region 5, was higher than its weld metal. The pit nucleation resistance values were in the range of 804.8–1014.2 mV and 611.1–747.8 mV for the welds WL and WH respectively indicating better passivation behavior exhibited by the joint which showed relatively smaller microstructural phase balance due to the reasons discussed above in the microstructure section. Overall, a general observation was that the regions with lower weld metal/base metal area ratio showed higher pit nucleation values than the regions with a higher weld metal/base metal area ratio. After EB welding, the pitting corrosion resistance order was observed to be: weld metal < composite zone (weld metal + heat affected zone + base metal) < base metal. This could be explained by the fact that in the unaged condition, the weld metal was the weaker zone, while base metal contained only the primary microstructural phases. The contribution of the base metal in the area fraction of the exposed surface suggested it to be beneficial for the overall corrosion behavior of the weld joints. Logically, it could be deduced that if the overall size of the fusion zone is small, EB welded joints as a whole could be expected to show better pitting resistance when subjected to corrosive environment in comparison to the joints with relatively larger size of the fusion zone with the fact that the weld zone may be the weaker one (in comparison to the base metal) in both the cases.

detrimental to pitting corrosion resistance. In the aged condition, region 3 of the weld WL showed small reactivation current density and the DOS was measured to be 1.3948 % which suggested a slight sensitization attributable to secondary precipitation including Cr and Mo-rich intermetallic sigma phase and carbides. Region 5 of the weld WH in the same conditions showed DOS of 0.01138 % which again indicated the absence of severe sensitization. The other regions of both the welds were found to be unsensitized. Thus, the overall assessment suggested that the welds were not severely sensitized after aging although secondary precipitation prevailed in them as indicated by the EDS and XRD results. This could be attributed to their low volume fraction [20] in the microstructure corresponding to the aging time used in the present work. Another contributing factor for not so significant vulnerability of these welds to IGC was the presence of Cr to an extent of 22 % which is higher as compared to the austenitic steels such as grade 304 which contains 18 % chromium. Thus, the localized depletion along the grain boundaries due to the presence of carbides is expected to be more prevalent in austenitic as compared to the duplex stainless steels. Moreover, the extent of depletion of vital alloying elements like Cr, Mo, and N along the grain boundaries is dependent on the volume fraction of the secondary precipitates and thus relatively harsher conditions may be required to render these welds susceptible to sensitization. For instance, Moura et al. [16] in their investigation on DSS S31803 involving the same test solution found that the DLEPR technique could not detect sensitization corresponding to an aged specimen with 4.7 % sigma phase. However, sensitization was detected corresponding to a relatively harsher aging condition that promoted 28 % sigma phase formation in the material. Furthermore, DOS values of less than 1 % (unsensitized condition) have been reported [22] for DSS S31803 in the same solution when aged at 650 °C for a duration of 316 min. These results from the literature suggest that sensitization could be detected in DSS corresponding to relatively longer aging time durations and for smaller durations, sensitization was not very significant. In a nutshell, it could be stated that for triggering intergranular corrosion in DSS EB welds, the volume fraction of secondary precipitates would play a pivotal role. Higher temperatures and prolonged exposures would possibly render these DSS EB welds vulnerable to sensitization which in the as-welded and corresponding to small duration aging conditions appeared to be largely unaffected. After the solution treatment of the aged welds, the reactivation current peaks were completely removed as observed from the reverse cathodic scans (Fig. 10c, f) suggesting the absence of sensitization. All the regions of both the welds investigated after solution treatment were found to be unsensitized. It could be deduced from this discussion that the EB welds were not susceptible to sensitization in the as-welded condition. Further, the aging condition of 850 °C/30 min also did not significantly sensitize the welds. However, harsher service conditions involving higher temperatures as well as longer exposure times may promote a larger volume fraction of secondary phases which in turn can make these welds prone to sensitization. Finally, the subsequent solution treatment of the EB welds led to the elimination of any reactivation current density peaks thus indicating the dissolution of deleterious phases that formed due to isothermal aging. 3.3.2. Potentiodynamic anodic polarization testing Potentiodynamic curves for different specimens are shown in Fig. 11(a–f) and the associated quantitative data are presented in Table 3. The base metal in the as-received condition showed pit nucleation resistance of 1154.3 mV which was highest amongst all the specimens in the unaged condition. Region 1 representing exclusively the weld metals of the specimens WL and WH showed the pitting resistance values equal to 804.8 mV and 611.1 mV respectively. The pitting resistance of the weld zones was relatively lower than the base metal. This could be attributed to rapid cooling in EBW that caused phase imbalance and probable nitride precipitation. Such precipitation 592

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Fig. 11. Potentiodynamic polarization curves for the specimens tested in 3.5 % NaCl solution at room temperature (a) BM and weld WL, T0 condition (b) BM and weld WL, T1 condition (c) BM and weld WL, T2 condition (d) BM and weld WH, T0 condition (e) BM and weld WH, T1 condition (f) BM and weld WH, T2 condition.

resistance were observed as compared to the as-welded condition due to aging-induced secondary precipitation. The weld metals of the joints WL and WH showed pitting resistance of 650.7 mV and 538.3 mV respectively indicating that the weld WL was superior to WH in the aged condition also. The overall evaluation showed pit nucleation resistance values of 400.8–723.7 mV and 443.2–681.2 mV for the welds taken in the same order. Interestingly, in the aged condition, the results of the region-wise examination did not show a similar trend of increasing pit nucleation resistance with decreasing weld metal/base metal area ratio and the joints as a whole showed lower resistance than their respective weld metals. This could be understood from the fact that in the aged condition, the base metal was also prone to secondary precipitation (unlike in the as-received condition), thus the weldments as a whole were rendered susceptible to corrosion attack which led to decrease in their pit nucleation resistance. It has been reported that the degree of sensitization (DOS) and pitting corrosion resistance follow largely an inverse relationship for AISI 304L stainless steel welds [23]. However, such type of trend was not observed in the present case. The aging condition which could adversely influence the pitting resistance of the DSS welds was found to be not so severe to affect their sensitization behavior. These observations suggest that relatively harsher conditions

Further, the examination of the specimens after testing divulged that both the base metal and EB weld metal suffered from pitting corrosion attack as shown in Fig. 12(a–d). The formation of stable corrosion pits could be observed on the surface of the as-received base material and the maximum pitting affected area was calculated to be 3801 μm2 for the pit represented by region 3 in Fig. 12b. The EB weld metal surface, on the other hand, showed not only different morphologies of pits but also underwent a larger surface degradation due to pitting as shown in Fig. 12(c–d). It contained independent pits as well as merged pits indicating that the microstructural phase imbalance that occurred in the weld zone was responsible for its relatively poor pitting behavior. For the weld zone, the maximum surface area damage due to pitting was calculated to be 7815 μm2 for the pit represented by region 4 as shown in Fig. 12c. Few larger pits in the weld zone that were formed due to the merger of neighboring pits had affected surface area as high as 14,338 μm2 for region 1 in Fig. 12c and 22,227 μm2 for region 7 in Fig. 12d respectively. Thermal aging at 850 °C/30 min degraded the pitting resistance of the base metal (Table 3) as compared to the as-received condition attributable to the formation of intermetallic σ-phase and carbides of type Cr23C6 and Cr7C3. For welds also, lower values of pit nucleation

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Fig. 12. Pitting corrosion on the surface of (a, b) base metal and (c, d) weld metal of DSS 2205 in 3.5 % NaCl solution along with calculated surface area affected due to pitting in both the cases.

4. Conclusions

in terms of aging temperature and time (which eventually decides the volume fraction of secondary phases) would be required to render these welds susceptible to sensitization but even short time aging would be able to affect their pitting corrosion property adversely when subjected to chloride containing environment. The solution treatment improved the degraded pitting resistance of the base metal due to the dissolution of secondary precipitation which was corroborated by its increased ferrite content. The region 1 comprising of the weld metal of the joint WL showed significant improvement along with a similar trend at all the regions tested. The joint as a whole (marked as region 4) showed greater pitting resistance than the weld metal. On the other hand, potentiodynamic polarization results for the weld WH also showed a shift of pit nucleation resistance to nobler values suggesting the pitting resistance restoration tendencies. Largely, all the regions showed escalated passivity with maximum values observed for region 5 representing the overall corrosion behavior of the joint. This could be explained based on microstructural phase balance improvement and the austenite formation which was promoted in both the welds, resulting in an overall enhancement of their pitting corrosion resistance.

The present work was focused on characterizing the intergranular and pitting corrosion behavior of electron beam welded 12 mm and 18 mm thick 2205 duplex stainless steel (DSS) joints using electrochemical techniques viz. DLEPR and Potentiodynamic anodic polarization respectively supported by assessment of their metallurgical characteristics. Further, the influence of post-weld aging and subsequent solution treatment on the aforementioned properties was also investigated. Following conclusions could be drawn based upon the results: 1 Fusion zones of both the weld joints were characterized by microstructural phase imbalance with ferrite proportion in excess as compared to the austenite. Relatively low welding heat input weld metal/fusion zone exhibited a relatively larger austenite proportion in the microstructure owing to larger grain boundary area facilitated by finer grain structure as compared to the high heat input weld metal. 2 Microhardness of both the weld metals was found to be higher than the parent metal. Thermal aging at 850 °C/30 min slightly decreased the microhardness but overall did not influence it with significant magnitudes. Subsequent solution annealing at 1050 °C/30 min imparted hardening to the welds attributable to restoration of their

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3

4

5

6

7

References

microstructural phase balance via partial dissolution of secondary precipitation caused by welding and post-weld aging treatment. Corrosion characteristics revealed that microstructural phase imbalance, owing to faster cooling rates experienced by the electron beam DSS welds, rendered them relatively more susceptible to pitting corrosion than the base metal, but the welds remained largely immune to intergranular corrosion (IGC). Further, the weld metal with a larger proportion of austenite showed relatively better pitting resistance than the weld with a larger microstructural phase imbalance. Thermal aging of the EB welds at 850 °C/30 min promoted the formation of Cr and Mo-rich intermetallic σ phase and carbides in the welds as indicated by XRD and EDS results, which further degraded their pitting corrosion resistance. DLEPR results of different zones of the welds revealed that they remained largely immune to sensitization under all conditions. Solution treatment of the welds at 1050 °C/30 min promoted the dissolution of intermetallic σ-phase and carbides back into the matrix. This led to the restoration of pitting resistance tendencies that were lost due to aging-induced secondary precipitation. Corrosion evaluation across different locations of the unaged welds, largely divulged that as the weld metal/base metal area ratio of the surface exposed to the test solution decreased, the pit nucleation resistance values were shifted to nobler values. In the as-welded condition, the pitting corrosion resistance order was observed to be: EB weld metal < composite zone (comprising of the weld metal, heat affected zone and base metal) < base metal. However, such an order did not prevail after thermal aging which could be ascribed to the enhanced vulnerability of both the weld metal and the base metal due to aging-induced secondary precipitation. This study on corrosion evaluation of the EB welded joints has established that, as fusion zones are more vulnerable to pitting corrosion as compared to their respective base metals, selecting welding procedures capable of resulting in small-sized fusion zones accompanied by lesser microstructural phase imbalance could impart better corrosion resisting characteristics to the welded structures meant for serving in corrosive environments.

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Declaration of Competing Interest The authors declare no conflicts of interest. Acknowledgements The authors are thankful to Er. Jaspal Singh Dhesi, Deputy General Manager-Welding Operations at Hendrickson, Pune (Maharashtra), India for the help extended during the experimental work. Further, the infrastructural support in terms of testing facilities provided by Welding Metallurgy Laboratory, Department of Mechanical Engineering, Sant Longowal Institute of Engineering & Technology (SLIET), Longowal, Sangrur-148106 (Punjab), India is gratefully acknowledged.

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