Contents lists available at ScienceDirect

Corrosion Science

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

Microscale investigation of the correlation between microstructure andgalvanic corrosion of low alloy steel A508 and its welded 309/308L stainlesssteel overlayer

Bright O. Okonkwoa,b, Hongliang Minga, Zhiming Zhanga, Jianqiu Wanga,⁎, Ehsan Rahimic,d,Saman Hosseinpoure, Ali Davoodic

a CAS Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, Chinab School of Materials Science and Engineering, University of Science and Technology of China, Shenyang 110016, ChinacMaterials and Metallurgical Engineering Department, Faculty of Engineering, Ferdowsi University of Mashhad, Mashhad 9177948974, IrandDepartment of Engineering and Architecture, University of Udine, Via Cotonificio 108, 33100 Udine, Italye Institute of Particle Technology (LFG), Friedrich-Alexander-Universität-Erlangen-Nürnberg (FAU), Cauerstraße 4, 91058 Erlangen, Germany

A R T I C L E I N F O

Keywords:Low alloy steelStainless steelSEMAFMInterfaceInclusion

A B S T R A C T

The microscopic scale correlation between microstructure and galvanic corrosion of low alloy steel-A508 and itswelded-309/308L stainless steel overlayer is investigated. Microstructural characterization shows that thewelding process induces complexities at HAZ of the base metal. The microstructure of HAZ changes fromtempered bainite to a non-uniform mixture of martensite and bainite, depending on the heat flow. This change inthe microstructure leads to a heterogeneous corrosion attack on the base metal (anode), while the weld metal(cathode) is protected in a galvanic interaction. The extent and rate of corrosion attacks on the surface con-stituents are quantified by AFM.

1. Introduction

Low alloy steel A508 autogenously welded with 309/308 L auste-nitic stainless steel cladding is a type of dissimilar metal weld (DMW)that is produced by shielded metal arc weld (SMAW) method. Suchjoint dissimilar metal structure possesses a combination of both con-stituent’s properties which can meet multiple objectives, especially inthe nuclear power plant (NPP) integral parts such as, reactor pressurevessel (RPV), RPV pipe nozzle, and steam generators. The RPVs ofboiling water reactors (BWR) and pressurized water reactors (PWR) areoften thick steel containers that hold nuclear fuel, when the reactors arein operation. Steel containers, thus provide the required barrier thatkeeps radioactive fission products away from the environment. Thesevessels and their pipe nozzles are often made of low alloy steel (LAS),which possess the required mechanical strength and is relatively cheap[1,2]. The inner walls, on the other hand, are protectively cladded withstainless steel (SS) which has a superior resistance to corrosion in ahigh-temperature environment [3]. The RPV which also acts as apressure barrier, plays a key role in the reactor’s safety.

The DMW which contains two or more dissimilar metals with dif-ferent chemical composition, offers a better-combined advantage by

improving the properties and usefulness of both metals [4]. The DMWjoining procedure is however worrisome because of the extreme con-ditions, such as weld thermal cycle and cooling rate, in which thewelding is performed. The welding of a base metal with a weld metalautogenously or with a buttering layer could metallurgically transformthe interface between the DMW and its surrounding areas [5,6]. Due tothe differences in the chemical composition, crystal structure, andthermal expansion of the base metal and weld metal, the microstructureof the fusion boundary (FB) and its neighbouring areas is often verycomplex [7–10]. These complexities can have a detrimental influenceon the mechanical and corrosion resistance of the DMW [11–14].

Recently, there has been the clamour to extend the operational li-cense of NPP’s beyond their anticipated lifespan (customarily 40 years),to about 50% longer than the initial period. This is due to the highcapital needed for constructing a NPP. In order for this to happen, therewould be additional demand for the reactor’s components to maintaintheir structural integrity for a time longer than expected. The prolongedextension in the service life of NPP reactors may introduce a variety ofpotential material aging issues that should be looked into in the NPP’slicense renewal process [15,16].

The main degradation issues associated with RPV steels and their

https://doi.org/10.1016/j.corsci.2019.03.027Received 18 December 2018; Received in revised form 25 February 2019; Accepted 16 March 2019

⁎ Corresponding author.E-mail address: wangjianqiu@imr.ac.cn (J. Wang).

Corrosion Science 154 (2019) 49–60

Available online 10 April 20190010-938X/ © 2019 Elsevier Ltd. All rights reserved.

T

nozzles are; embrittlement caused by radiation-induced solute defectscluster hardening, corrosion, and stress corrosion cracking (SCC) phe-nomena [15,17]. Most corrosion problems in NPP caulescent from poorwater chemistry control. Accidental ingress of sea water which containsharmful ions (such as Cl− and SO4

2-) from the condenser into the RPVhas been reported in Hamaka nuclear power plant in Japan [10,18].Also, the hazardous effect of leakage of coolant into the head of RPV hasbeen reported in Davis – Besse reactor PWR in USA [15]. Such eventscan give rise to pitting [19] and SCC [20] damages in stainless steel,thereby exposing the low alloy steel, which in turn give way for gal-vanic corrosion to set in between the dissimilar metals. Furthermore, ithas been reported that the stainless steel cladding could be mechani-cally damaged when the penetration component is being replacedduring repairs [21]. Another possible mechanical damage occurs whenthe SS cladding becomes thinned due to wear of the dropped objects,such as cracked bolts and remnant tools. In the worst case, base metalmay be exposed and thinned, also by wear and galvanic corrosion.Therefore, to consider the reliability of the pressure vessel, it is neces-sary to study the galvanic corrosion that may occur as a result of theexposure of the base metal to a corrosive environment. This is im-portant for scientists and engineers in this field, in order to enable themto provide useful data for risk assessment models necessary for main-tenance and service performance of the NPPs components (especiallythe RPV).

Several studies have been conducted on the galvanic interactions ofDMW’s [22,23]. Wang S. et al. [23], studied the characterization of lowalloy ferritic steel – Ni base alloy dissimilar weld interface and observedthat the interface region with a wider spread of martensite layer had alower Volta potential compared to the region with a thin layer. Theyconcluded that galvanic corrosion was the prevalent corrosion me-chanism of the A508 - alloy 52 interface region. Weng S. et al. [24], alsoobserved that the controlling mechanisms in corrosion of Ni-Cr-Mo-Vsteel welded joints in 3.5 wt% NaCl solution (at 180 °C and 0.8MPacondition) were galvanic corrosion and stress assisted corrosion.However, the coarse grain heat affected zone (HAZ) was more suscep-tible to corrosion. This was attributed to chemical composition andmicrostructure of this region. Based on the studies which are carried outso far, it is inferred that galvanic corrosion of DMW is controlled byseveral factors including; microstructure and chemical composition[25–28], temperature [29], stress [30,31], and the area ratio of anodeto cathode [32,33].

The microstructure of SA508–309/308L DMW part of the dissimilarmetal weld joint (DMWJ) (schematically shown in Fig. 1a) has beenextensively studied by our research group and other researchers[34–36]. For instance, Ming H. et al. [34], investigated the micro-structural characterization of SA508–309/308L–316 DMW safe-endjoint, and observed the appearance of HAZ with a complex structure inLAS near the fusion boundary (FB). They found out that, as a result ofcarbon migration from the LAS to SS, the hardness of LAS decreased atsome zones very close to the FB (carbon depletion zone), and carbon-rich zones appeared in the SS part of the interface. This phenomenonbroadened the difference in properties within the fusion zone. Xiong Q.et al. [36], categorized the HAZ of LAS 508 III into three zones whichare: intercritically reheated coarse grain zone, fine grain zone, andcoarse grain zone. In some cases, chromium enriched precipitates wereobserved in the 309 L austenitic stainless steel part, and some authorslinked this observation to tertiary phase (sigma phase) formation due tothe post weld heat treatment (PHWT) and material aging [37,38]. Thisphase has been pointed out to be the main cause of welded SS de-gradation [39]. However, the review indicates that little or no work hasbeen done so far in relating the microstructure of SA508–309/308LDMW to the galvanic corrosion occurring between dissimilar parts on amicroscopic scale.

The aim of this work is to improve the understanding of the galvaniccorrosion damage mechanism between A508–309/308L SS dissimilarweld coupled metals in PWR primary water simulated solution

containing sulphate impurities. This ultimately would provide valuableinformation on the corroding surface constituents, and corrosion phe-nomena occurring between the galvanic couple. We used optical mi-croscope (OM) and scanning electron microscope (SEM) to characterizethe microstructure of the DMW. We also employed scanning vibratingelectrode technique (SVET) and atomic force microscope (AFM) tech-niques to study the corrosion processes occurring on a microscopicscale. Detailed analyses of the AFM results, using histograms obtainedby the multimodal Gaussian distribution fit, and power spectral density(PSD) analysis [4,40], enables prediction of corrosion rates of corrodingalloy constituents and provides new insights on the role of differentcorroding surface constituents in the global degradation of the DMW.

2. Experimental

2.1. Materials

The test samples (A508-309/308L SS) used in this work were cutfrom a domestic DMWJ, as shown in Fig. 1b. Obtaining a small cutsample of the RPV is quite difficult; therefore, the RPV pipe nozzle wasused instead of a RPV. Both RPV and RPV pipe nozzle, however, sharethe same joining technique of the A508-309/308L SS dissimilar metalweld, which has been reported previously [1,23,41]. Table 1 shows thechemical compositions of the base metal and combined stainless steelweld cladding. The sample preparation procedures are expatiated in the

Fig. 1. (a) Schematic of a dissimilar metal weld joint safe-end. (b) Test samplelocation.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

50

introduction of each experimental section.

2.2. Microstructure analysis

Metallographic examination of the test samples was carried outusing OM and SEM. The OM specimen was cut in a dimension of15mm×15mm×3mm. Prior to the microstructure observation withan optical microscope (Carl Zeiss Axio Observer. Z1m), the sample wasground with SiC abrasive paper sequentially with a grit size rangingfrom 150# to 2000#, and mechanically wet polished with 2.5 μmdiamond paste. To enable clear visual of the constituent phases andgrain boundaries, the LAS A508 was etched with 4% nital solution(97% ethanol + 3% nitric acid). The stainless steel 309/308L waselectro-etched in 10% oxalic acid + 90% H2O solution for 10 s, with adirect current of 5 V. These two different etchants, which are the knownprocedures [34,36] were used based on the different characteristics andetchant response of the base metal and weld metal.

Scanning electron microscope equipped with energy dispersive x-ray spectroscopy device (SEM/EDX, FEI XL30), was used to obtainhigher magnifications of the microstructure (compared to OM) andchemical composition across the interface region and selected areas inboth the base metal and weld metal. Same sample and preparationtechniques as that of the OM sample were used for the SEM evaluations.

2.3. Scanning vibrating electrode technique (SVET)

SVET is an electrochemical method that provides a naturally oc-curring way of defining the intensity and process of corrosion attacks,taking place in a metal submerged in an electrolyte. SVET maps areconstructed by showing the current density maps of the anodic andcathodic activities happening above the corroding surface [42,43]. Forthe SVET test sample, a region of SA508-309/308L DMW which in-cluded the interface was cut in a dimension of 4mm×4mm×3mm,of which the A508 part was 2.1mm×4mm×3mm and the 309/308 L SS cladding was 1.9mm×4mm×3mm. The sample piece wasfixed vertically in epoxy and was mechanically ground and polishedwith SiC paper of various grit sizes (150# to 2000#) and diamondpaste, respectively. The electrochemical cell was set up by sticking alayer of sellotape over the epoxy containing the sample, in order to keepthe solution on the metal surface throughout the experiment.

The SVET experiments were performed using commercial SVETdevice (Applicable Electronic Inc., USA) controlled by ASET 2.0 soft-ware (Science Ware Inc. USA) for 28 h, at an interval of 3 h. PWR pri-mary water simulated solution containing 1500 ppm B as H3BO3 +2.3 ppm Li as LiOH was used as the electrolyte. 96 ppm sulphate(Na2SO4) impurity was added to the PWR primary water simulatedsolution, because of the inability to obtain a good current density valuewith the PWR primary water simulated solution, owing to its very lowconductivity (of 33 μs/cm). Also, it has been observed that there aredeposits of condensed corrosion products at the bottom of the RPV,which may act as impurities in the RPV contained primary water. Thevibrating electrode used was an insulated Pt/Ir probe, which had pla-tinum black deposited on its 15 μm diameter spherical tip. The sampledistance from the probe was 100 μm, and the probe was vibrating at anamplitude of 30 μm in a plane perpendicular (z) and parallel (x) to thesample’s surface. The vibrating frequencies were 270 and 870 Hz, re-spectively. Each SVET scan took ˜20min to complete.

2.4. Immersion test and atomic force microscopy (AFM) analyses

The AFM test sample was cut from the same interface spot as that ofthe SVET sample, in the dimension of 10mm×10mm×2mm. Thesample was then immersed in PWR primary water simulated solution asstated above, but with a different concentration of sulphate impurity(2880 ppm SO4 as Na2SO4) for 6 h, at ambient temperature in air. Thesulphate concentration of the electrolyte was different in the SVET andAFM tests, as lower electrolyte conductivity was required in the formerto obtain a gradual corrosion phenomenon. However, the sulphateconcentration used for both tests was still in the range contained inseawater [44]; thus simulating a corrosion process similar to an acci-dental ingress of seawater from the condenser into the RPV (flowingthrough the RPV pipe nozzle DMWJ). Surface topographies of the sameinterface area were observed and analysed with AFM, before and afterimmersion, using the commercial Solver-Next AFM software (NT-MDTNova). This was achieved by making two micro-indentation marks atboth sides of the DMW interface region confining the mapped area. Thesurface topography mapping was carried out using the contact modetechnique, at a scan rate of 0.5 Hz and a pixel resolution of 512× 512.To better interpret the topography data, a combination of multimodalGaussian distribution based on histogram and power spectral densityanalysis were performed. PSD analysis, based on fast Fourier transform(FFT) of AFM images, simultaneously provides both lateral and long-itudinal information about the topography distribution of surface con-stituents at different corroding regions, as compared to other statisticaltreatments of AFM images, such as root-mean-square, mean value,surface skewness, and peak-to-peak analysis. The PSD analysis thusprovides a more quantitative evaluation of the results, as compared tothe conventionally used line scan analysis. A correlation between in-dividual surface constituents of the microstructure and estimated localcurrent densities were also recognized.

3. Results and discussion

3.1. Microstructure characterization

3.1.1. OM examinationFig. 2a displays the microstructure of the A508-309/308L SS close

to the fusion boundary. The low alloy steel has a tempered bainitemicrostructure, which consists of ferrites and minuscule carbide net-work formed along the grain boundaries of the ferrite. Traversing to-wards the fusion boundary, the microstructure of A508 base metalbecomes complex and non-uniform. The 309/308L weld metal consistsof a worm-like ferrite (dark grey δ-phase) distributed randomly in theaustenite matrix (light grey ϒ-phase). However, an irregular brightplain area of solely austenite phase is also observed in the SS part ad-jacent to the fusion boundary.

Theoretically, the effect of heat dissipation (from the welding andother pre or post-weld heating conditions) toward the base metal de-creases with increasing distance from the fusion boundary. This leads tothe typical four clear-cut regions in the HAZ, namely; coarse-grain re-gion reheated above 1100 °C, fine-grain region reheated within therange of 900 °C – 1100 °C, partially refined region reheated within700 °C – 900 °C, and tempered region reheated below 700 °C [34].Nevertheless, the HAZ microstructure may vary slightly from this ca-tegorization when the multi-pass and multi-layer welding process is

Table 1The chemical composition of the base metal and the weld metals (wt.%).

Alloy Cr Ni Mn Si C Mo Cu S Nb Fe

SA508 III 0.16 0.82 1.40 0.23 0.192 0.491 0.039 0.0007 <0.001 Bal.309L 23.94 13.48 1.67 0.40 0.016 0.031 0.044 < 0.001 0.001 Bal.308L 20.38 10.32 1.76 0.37 0.020 0.022 0.029 0.002 0.001 Bal.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

51

performed. The trend of microstructure evolution from the FB to theA508 matrix (shown in Fig. 2b) are; coarse grain refined region (mar-tensite+ little deposit of gritty bainite (Fig. 2bi and bii)) → fine grainrefined region (fine martensite+ little deposit of fine bainite (Fig. 2biiiand biv)) → partially refined region (bainite+ little deposit of finemartensite (Fig. 2bv)) → tempered region (bainite (Fig. 2bvi)). It isevident from the OM image, that the microstructure of the coarse grainrefined (CGR) region (Fig. 2c) is totally distinct from other regions, likethe fine grain refined (FGR) region (Fig. 2d). These variations in themicrostructure across the fusion boundary are due to high heat influx,which leads to different recrystallization of the grains depending ontheir distance from the fusion boundary.

The transition of microstructural phases from the base metal to the

weld metal can be validated using the Schaeffler steel prediction dia-gram [37,45] presented in Fig. 3. In this diagram, the chromiumequivalent (Creq) values are represented on the x-axis, while the nickelequivalent (Nieq) values are represented on the y-axis. The Creq(Creq= Cr%+Mo%+1.5Si%+0.5Nb%) and Nieq (Nieq=Ni% +30C%+0.5Mn%) obtained from the chemical composition of thevarious metals shown in Table 1, are 0.9 and 7.5 respectively for A508,and 24.6 and 14.8 respectively for 309L SS. These values were input inthe Schaeffler diagram which showed that the A508 steel can be foundvery close to the ferrite+martensite phase region, while the 309L SS isat the austenite+ ferrite region, thereby confirming the OM results.The plot in Fig. 3 also confirms the possibility of having a martensiticlayer, being formed along the FB of A508-309/308 L welded joint,

Fig. 2. (a) Optical micrograph image of A508 – 309/308L SS dissimilar metal weld. (bi – bvi) Magnified images of the blue rectangular box area in (a), showing themicrostructure transtion in A508 HAZ: martensite+little deposit of gritty bainite (CGR) → fine martensite+little deposit of fine bainite (FGR) → bainite+ littledeposit of fine martensite (PRR) → bainite (TR). (c) and (d) Higher magnification images of CGR and FGR regions respectively (For interpretation of the references tocolour in this figure legend, the reader is referred to the web version of this article).

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

52

depending on the dilution ratio, in agreement with the SEM results. It isobserved from the Schaeffler diagram that during solidification, theweld metal transverses slightly through the sole austenite phase (i.e atthe point it moves away from the FB) before moving further towards theaustenite+ ferrite (A+F) phase, just as seen in the OM result. Thereason for this solidification behaviour in 309/308L SS is that an in-crease in carbon content in the 309/308L interface region, due tocarbon diffusion from the base metal into the weld metal duringwelding [46], leads to a corresponding increase in the Nieq. A higherNieq in the Creq/ Nieq, in turn, favours austenite phase formation andvice versa for the ferrite phase.

3.1.2. SEM examinationSEM examinations are carried out before and after etching of the

test sample. Fig. 4a shows a typical inclusion present in A508, whosechemical composition was determined based on the EDS spectra shownin Fig. 4b. These inclusions, which are randomly distributed in differentshapes and sizes in A508, contain elements of O, Al, S, Mn, and Fe. Thisdistribution of elements suggests that MnS and oxides of Al in the in-clusion are nucleated on an oxide particle during the solidification ofA508 [22]. The elemental composition of the EDS spectra showed thatMnS was prevalent in the inclusion, while the high peak of Fe ischaracteristic for the steel matrix.

SEM micrograph of etched A508-309/308L SS interface region ispresented in Fig. 5a. This image clearly shows the lath martensiteformed on the coarse-grain refined region of A508, and at the fusionboundary. This martensitic transformation occurs as a result of heatflow, cooling rate, and interfacial migration of elements between thetwo metals during the welding process. It has been shown that Ni mi-gration from the weld metal into the base metal results in sub-grain

refinement, which leads to an increased fraction of martensite in theHAZ [47]. Six points were selected across the interfacial region andtheir chemical compositions (obtained from their EDS spectra) areprovided in Fig. 5b. This histogram plot presents the migration trend ofFe from the base metal to the weld metal, and vice versa for the Cr andNi migration trends. Point 1 which is the farthest point away from theFB, contained mainly Fe and little amount of Ni and Cr, respectively.Moving close to the FB from the base metal matrix (points 2 and 3), theCr and Ni content increased, while the Fe content decreased. A reversetrend is observed moving away from the FB into the weld metal (points4–6). This trend in the change of chemical composition across the FBconfirms the existence of an element migration phenomenon occurringalong the interface region, during the welding process. This phenom-enon leads to increased heterogeneity within the interface [1,46],which may be detrimental to the DMW.

Fig. 6a and b show the SEM micrograph of the various phases pre-sent in SS weld cladding. Inserted at the edges of both images are theEDS spectra of the marked regions, which show peaks of Fe, Cr, and Niin the primary phases (austenite, ϒ, and ferrite, δ), and tertiary phase(Fe-Cr-Ni ternary sigma, σ). From the EDS spectral analysis and che-mical compositions shown in Table 2, it is evident that the ternarysigma phases have the highest chromium content, followed by theferrite and austenite phases, respectively. The chromium-rich sigmaphase often precipitates as a result of multi-pass welding [48] and/orpost weld heat treatment (PWHT) of the A508-309 L DMW. The PWHTat a temperature range of 760 °C – 815 °C on DMW’s [49] aims at re-ducing the residual stress at the interface. However, it may cause theprecipitation of the deleterious sigma phase at the temperature range of600 °C – 800 °C [37] in SS containing above 20% chromium content.The sigma phases in Fig. 6a are nucleated at the ϒ/δ phases interface(σ1), and at the bulk δ phase (σ2). The reasons for such distribution ofsigma phases are: (i) the ϒ/δ phases interface is a high energyboundary, which makes it suitable for diverse nucleation [37], and (ii)the δ phase transforms easily at elevated temperature into the σ phasedue to its high chromium content and BCC crystal structure, which areloosely packed and accommodates faster diffusion [50]. The formationof these chromium rich σ phases may indirectly lead to chromium de-pleted regions, which could adversely affect the corrosion resistance ofthe material.

Referring back to the Schaeffler’s steel prediction diagram in Fig. 3a,it is noticeable that the 309 L of SS is found in the embrittlement region(brown shaded region, also known as the sigma phase region). Thisconfirms the possibility of σ phase precipitation at the 309 L of thestainless steel, because of its higher chromium content.

3.2. Galvanic corrosion behaviour of the coupled A508-309/308Ldissimilar metal weld

Fig. 7a–j present the SVET maps for the A508-309/308L SS coupled

Fig. 3. Schaeffler diagram containing the embrittlement region of σ phase, usedfor predicting the microstructure of steels [37,45].

Fig. 4. (a) SEM image of A508 before etching. (b) EDS spectra and elemental composition of the inclusion in A508 (a) base metal.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

53

dissimilar metal immersed in the solution (stated above in section 2.3)as a function of immersion time. In all SVET maps, the A508 base metalis situated on the left, while the 309/308L weld metal is situated on theright. The colour bar at the side of each map represents the anodic(positive) current density distribution, and the cathodic (negative)current density distributions obtained for each scan. The former is re-lative to the flux of cations detected by the SVET probe from the surfaceto the solution; while the latter is relative to the flux of anions detectedas well by the SVET probe from the surface to the solution. After 3 h ofintroducing the coupled dissimilar metal in the PWR primary watersimulated solution containing sulphate impurities (with a conductivityof 80μs/cm), ionic dissolution is observed on A508 base metal as shownin Fig. 7a. Interestingly, high anodic activities are observed at the CGRregion adjacent to the fusion boundary in A508; which included loca-lized anodic current density, suspected to be pitting corrosion induced

by the dissolution of MnS inclusion. Cathodic current is observed tocover a larger part of the other regions, owing to the magnetite oxidefilm growth on the base metal, which retards corrosion. This lateraldistribution of anodic and cathodic sites suggests an uneven anodicdissolution, which results in a heterogeneous attack on the A508 basemetal. For the 309/308 L SS, a total cathodic current is detected overthe whole surface, owing to the protective chromia oxide film formed.

In the course of sample immersion to the corrosive medium, themetal being dissolved in A508 is Fe. However, there are still pockets ofdissolutions of Mn from the inclusions present in the base metal[22,51]. These metallic dissolutions drive the oxidation reactions oc-curring on the base metal surface, as shown in Eqs. (1) and (2), re-spectively. Based on the SVET maps, the 309/308 L SS weld metalsurface shows only cathodic currents up to 28 h immersion. It is thussuggested that the dissolution of Cr from the weld metal leads to theformation of chromia protective oxide film, which retards further ionicdissolution. It has been reported previously by Niu, L. et al., [52] that instainless steel, Cr preferentially reacts and oxidizes in the presence ofFe, due to its higher tendency to oxidization (standard electrode po-tential of Cr and Fe are −0.74 V and −0.44 V, respectively). It istherefore suggested that the oxidative reaction shown in Eq. (3), takesplace in 309/308L SS before the formation of the passive oxide film.Despite the possibility that the dissolution of metallic ions initiallyoccurred in both metals, it was, however, more prominent in A508 dueto its lower corrosion potential than that of the 309/308L stainlesssteel.

→ ++ −Fe Fe 2e2 (1)

Fig. 5. (a) SEM image of partly etched A508 – 309/308L SS interface region. (b) The histogram plot of the elemental composition obtained from the SEM-EDS areascan across the interface.

Fig. 6. (a) and (b) SEM image and EDS spectra of the austenite(ϒ), ferrite(δ) and sigma(σ) phases in 309/308L Stainless steel weld metal.

Table 2The chemical composition of the different phases in 309/308L stainless steelweld cladding.

Phase Cr (wt.%) Fe (wt.%) Ni (wt.%)

ϒ1 18.73 68.62 12.65ϒ2 19.46 67.98 12.56δ1 23.92 67.68 08.40δ2 21.93 67.37 10.70σ1 27.27 66.96 05.77σ2 29.69 65.74 04.57

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

54

Fig. 7. (a–j) SVET maps of A508 – 309/308L SS DMW couple, immersed in PWR primary water environment (1500 ppm B as H3BO3 + 2.3 ppm Li as LiOH)containing sulphate impurity (96 ppm SO4 as Na2SO4) for 3 h, 6 h, 9 h, 12 h, 15 h, 18 h, 21 h, 24 h, 26 h and 28 h respectively. (k) Optical microscope image of theSVET sample after immersion in the corrosive media. (l) A plot of average current density with time, extrapolated from same data points within the black rectangularbox area in (a), for all scans.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

55

+ → + + ++ − + −MnS 3H O 2Mn S O 6H 8e22

2 3 (2)

→ ++ −Cr Cr 3e3 (3)

When the cations that are being detected as positive current den-sities (anodic current) dissolve, a corresponding cathodic (reduction)reaction occurs in the electrolyte, which balances up the electro-chemical process. Since the pH of the electrolyte is slightly acidic (˜pH6.2), the cathodic reactions occurring in this system are predominatelyoxygen reduction and/or water reduction (reactions in Eqs. (4) and (5))[53]. To account for the changes in the water content of the electrolyte,which was observed during the experiment from time to time pre-sumably due to evaporation, pipette drops of the electrolyte wereconstantly added to maintain a fixed volume.

+ + →− −O 2H O 4e 4OH2 2 (4)

+ → +− −2H O 2e H 2OH2 2 (5)

As stated earlier, the oxidation and reduction reactions occurring onthe base metal and weld metal lead to the formation of oxide films onthe surfaces of both metals. While magnetite (Fe3O4) is graduallyformed on A508 [52] through an active-passive mechanism, chromia(Cr2O3) is rapidly formed on 309/308L SS [21] through a spontaneous-passive mechanism. Eqs. (6) and (7) show the reactions leading to theformation of these oxide films. In spite of the fact that both oxide filmsprovide protective barriers to both metals, the chromia formed on SS ismore compact, stable and adhere strongly to the metal substrate [54].The chromium oxide thus gives SS its enhanced corrosion resistiveproperties.

+ → → + ++ −3Fe 6OH 3Fe(OH) Fe O 2H O H22 3 4 2 2 (6)

+ → → ++ −2Cr 6OH 2Cr(OH) Cr O 3H O33 2 3 2 (7)

As the test progresses, the anodic current in A508 decreases (asshown in Fig. 7b and c) until it reached a state where only totalcathodic current density is detected on the whole coupled dissimilarmetal surface. This is observed at the 12th h of immersion (within thestudied time range) (Fig. 7d) and lasted till the 15th h of immersion(Fig. 7e), owing to the total passivation caused by the various oxidefilms formed on the base metal and weld metal (magnetite and chromia,respectively). Uyime D. et al., [55] reported that SVET records a totalnegative current value when the SVET probe detects no anodic current.Other authors [56,57] have reported a similar phenomenon. The reasonfor this is that the SVET probe which can detect ions present on thesurface of the metal, only detected OH−, O2- and other negative ions(such as SO4

2-) mainly present in the passive film, and at the interfacebetween the oxide film/electrolyte during this period, owing to theprobe to sample distance. Although, there could be possibility of slightcorrosion occurring at the metal/oxide film interface of the A508, but atthis time could not be detected by the SVET probe due to the protectivenature of the magnetite oxide film.

Fig. 7f displays the SVET map after immersion for 18 h. Interest-ingly, at this time the passivation in the A508 ruptures at the CGR re-gion adjacent to the FB, and small anodic current density values sig-nifying corrosion are detected. Although the magnetite oxide filmgrowth rate decreases (with obvious lower cathodic current density), itis still noticeable in other areas of the A508. This reveals the highsusceptibility of the CGR region to corrosion. The passivity breakdownat the oxide film/electrolyte interface, in this case, might have involveddefects in the passive film [19], a phenomenon which required more indepth studies.

With increasing the immersion time the corrosion phenomena oc-curring in the A508 base metal, begin to spread over the surface untilthe whole surface is covered. This is evidently shown in Fig. 7g–j. Thestainless steel is observed to be cathodic throughout the duration of thetest. This reveals the protective nature of the chromia oxide film andconfirms the galvanic interaction between the two metals, in which thelow alloy steel corrodes in detriment to the stainless steel.

After immersion for 28 h, the test sample was carefully retrievedfrom the SVET sample holder, air dried and observed with an opticalmicroscope. Fig. 7k shows the corrosion products that are formed onthe A508 surface while the 309/308L stainless steel surface is almostunharmed. Also from the image in Fig. 7k, it is evident that the CGRregion adjacent to the FB in A508 is the most corroded area, with amore brownish-yellowish corrosion product. This confirms the bene-ficial effect of the galvanic interaction between the coupled dissimilarmetal weld on stainless steel.

Fig. 7l displays the plot of average current density with time, ex-trapolated from the same data points within the black rectangular boxarea in Fig. 7a, for all scans. These areas are the most susceptible re-gions to corrosion (CGR region in A508), owing to the presence ofdifferent high active phase constituents in this region, as pointed outearlier in the microstructural characterization. From the plot, the sug-gested mechanism of corrosion is categorized as follow; start-up ofcorrosion (3 h) → to decrease in corrosion as a result of passivationbuild up (6 h – 9 h) → total passivation (oxide film growth (12 h)) →decrease in oxide film and eventually, rupture of the oxide film(15 h – 21 h) → general corrosion increase (24 h – 28 h). To further un-derstand the reason for the higher susceptibility of the CGR regionadjacent to the FB to corrosion, AFM analyses are conducted which arediscussed in the next section.

3.3. AFM characterization of the corrosion attack on the surfaceconstituents

Fig. 8 shows the topography of the test sample before and after c.a.6 h immersion in corrosive solution (as stated in section 2.4). While309/308L SS (right side of the figures) seems totally unattacked, A508base metal (left side of the figures) clearly undergoes aggressive gal-vanic attacks at the weld borderline region. In the presence of this

Fig. 8. AFM Topography of A508 – 309/308L SS DMW couple before, (a), and after, (b), immersion for 6 h, in PWR primary water environment (1500 ppm B asH3BO3 + 2.3 ppm Li as LiOH) containing sulphate impurity (2880 ppm SO4 as Na2SO4).

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

56

galvanic couple, passivity enhancement by passive layer growth is ex-pected on the 309/308 SS side. However, the growth of the passivelayer is in the order of a few nanometers, which is undetectable in thisscale of imaging. Therefore, the effect of such passive layer growth isnot considered in the calculation and subsequent quantitative assess-ment of the topography data. As a result of alloy mixing (discussed inSection 3.1.2), a slight nobleness is observed exactly at the weld bor-derline particularly toward the 309/308L SS. More corrosion attack isobserved on the newly formed martensite phase, whereas the temperedbainite matrix is less affected. Minuscule carbide network that is visibleon the grain boundaries of tempered bainite phases, as reported pre-viously [36] is barely attacked compared to the matrix. The Minusculecarbide network appears as a more noble phase because of its content ofchromium, similar to what has been reported in previous Volta poten-tial mapping [58,59]. The most prone surface constituents to corrosionare detected to be the inclusions, mostly comprising of MnS [47,60], onA508 alloy. MnS dissolution and consequently pitting corrosion occurssimultaneously with galvanic corrosion in the weld borderline region,with the MnS dissolution being the dominant corrosion mechanism ofcorrosion damage. This confirms the high localized anodic currentdensity observed in the CGR region of the base metal adjacent to the FBin Section 3.2.

In order to perform a deeper assessment of the extent of attacks onvarious surface constituents, histogram and PSD analysis of AFMimages are performed. To interpret the topography histograms, multi-modal Gaussian distributions are calculated using Eq. (8) [40,61]:

= ⎡⎣⎢

−− ⎤

⎦⎥Y 1

σexp

2(x μ)σπ

2

2

2(8)

where Y represents the count’s number, σ is the standard deviation, x isthe topography value and μ is the mean height value of AFM topo-graphy. A meaningful correlation can be found between height valuesas roughness component with respect to their spatial frequency vector,which can be explained by the following equation [40,61]:

∫= −→∞

−

PSD(f) lim 1A

z(r)exp( 2πif.r)drA

A

2

(9)

where z(r) represents the height data of the surface roughness, A is thesurface area of the measuring field, r is position vector and f is thespatial frequency vector in the −x y plane.

Fig. 9 shows the PSD and multi-modal Gaussian histograms of to-pography results before and after immersion to corrosive media. Ac-cording to PSD analysis, the lowest and the highest topography((height)2 and (x–y)2) values on PSD profiles are related to the highestand the lowest spatial frequencies, respectively. Additionally, differentslopes are observed in the PSD curves, representing a heterogonousroughness distribution on the surface. According to PSD profile and theheight value of components in Fig. 10, MnS inclusions exist in the highfrequencies while the lower frequency values are related to martensiteregion followed by the bainite regions, before 1.07 μm−1. The regionsaround the fusion boundary (represented by dashed line in Fig. 9a)show uneven changes of topography distribution, which is related tothe complex microstructure formation at the weld interface toward309/308 L SS. As shown in Fig. 9a, the PSD profile of the sample afterimmersion test represents the higher PSD distribution values than thatbefore immersion test in the entire spatial frequencies. Furthermore,the slope of the PSD profile in the frequency values larger than 1.07 μm-

1 decreases after the immersion. These observations are due to the moresevere corrosion attacks on A508 base metal and slight passive filmformation on 309/308 L SS regions, in accordance with the SVET resultsand microscopic evaluations.

According to the histogram graphs in Fig. 9b and c, two and fivemodals are identified in AFM images before and after immersion, re-spectively. Each modal can be directly correlated to the height values of

the surface constituents of the base metal and weld metal. In Fig. 9b,the little variation in height between the base metal and weld metal isthe result of mechanical polishing imperfections, which invariablycannot be totally eradicated. After immersion, new modals appear inthe histograms. The five modals in Fig. 9c represents the surface cor-rosion height difference in the following decreasing order; 309/308LSS > weld borderline > tempered bainite in A508> tempered mar-tensite in A508>MnS inclusion in A508, as summarized in Table 3.Existence of these modals after immersion of the sample to corrosivemedium clearly indicates localized corrosion attacks on the DMW.Fig. 10 shows the color-map of AFM image which is used for finding thecorrelation between the surface constituents and topography feature.This image is used to extract the Gaussian distribution parameters in-cluding the average and standard deviation of heights before and aftercorrosion attack on individual surface constituents.

The corrosion rate of individual surface constituents is estimatedbased on the dissolution rate of three dissolved phases, i.e. martensitephase, tempered bainite matrix, and MnS inclusions and the calculatedcorroded depth and corrosion rates are provided in Table 4. The cor-rosion rates are calculated using Faraday law (Eq. (10)) [62];

×=

ρ.A. d nFM. t

I (10)

where ρ is density, A is the surface area, d is reduced thickness, n isthe number of transferred electron, F is Faraday constant, M is mole-cular weight, t is the time of exposure, and I is current. (Molecularweight of 57.5 g/M and 87 g/M are used for LAS 508 and MnS, re-spectively, with the density of MnS being 3.99 g/cm3).

As can be seen in this table, the electrochemical dissolution rate ofthe MnS inclusion in LAS 508 is the fastest occurring reaction (re-presented in eq. 2) followed by dissolution of martensite phase andtempered bainite matrix, respectively.

As seen in Table 4, despite that MnS has a very small surface area,its dissolution induces a corrosion current density of more than 130(μA cm−2), which is c.a. 4–6 times higher than that in the martensitephase or tempered bainite matrix. This is an interesting issue since instainless steel, the remained surface as a passive layer is believed to beinert (compared to the MnS pits). Here on LAS 508 alloy, in addition tothe high current density corresponding to MnS, a moderate dissolutionof the alloy matrix is also observed which is not negligible.

The overall corrosion rate of the DMW, as well as the corrosion ratesof its individual constituents, can be compared based on SVET currentmappings and AFM based corrosion rate estimation. The slightly dif-ferent corrosion rates estimated from the two techniques are related tothe difference in sulphate concentration of the electrolytes and thelimitation of SVET in accurately determining the current densities.These limitations include; the movement of the SVET probe and itsvibration which sometimes leads to more transfer of oxygen to surface,and unbalanced reactions between anodic and cathodic currents onsame region [63] In addition, the sensitivity of SVET technique andcalibration condition must be controlled during mapping in order toderive the absolute corrosion rates. This approach of estimating cor-rosion rates based on the AFM results, provides further insights on thegalvanic interaction of A508–309/308 L SS DMW in a corrosivemedium, and it becomes evident that high corrosion attack rates of theMnS and martensite phases are responsible for the high susceptibility ofthe CGR region adjacent to the FB to corrosion.

To evaluate the corrosion attack predicted by the AFM data, the testsample after immersion for 6 h in corrosive medium, was cleaned withethanol, blow dried and observed with the SEM. The SEM image pre-sented in Fig. 11 confirms the presence of pits randomly distributed inthe CGR region of the A508 base metal. The image showed clearly thecorrosion attack on the A508 alloy and the almost unattacked 309/308 L stainless steel surface. The box inserted in the picture gives theelemental composition of the pit at the area marked with a cross. Theelemental composition revealed that the pit still contained a remnant of

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

57

Mn and high percentage of Fe, confirming that this was indeed an in-clusion site. Therefore, SVET, AFM, and SEM observations on the DMWsample after immersion to the corrosive medium provide a unified

conclusion that pitting initiates at the inclusion sites in A508 alloy andthat galvanic interaction occurs between the A508-309/308 L DMWcouple.

4. Conclusion

The correlation between microstructure and galvanic corrosion ofA508–309/308 L SS weldment was studied in microscale using OM,SEM, SVET, immersion, and AFM techniques. From results obtained,the following conclusions can be inferred:

1) Microstructural investigation clearly revealed the complexity in themicrostructure of the HAZ in A508 base metal, owing to the weldingwith stainless steel and PWHT process. The tempered bainite mi-crostructure of the base metal matrix changed tremendously movingcloser to the fusion boundary of the joint, into a complex mixture ofthe martensitic and bainitic structure. The element dilution zone atthe interface was observed to be slightly different from both metalmatrixes. This led to enhanced property difference within the in-terface. Primary (ϒ, δ) and chromium-rich tertiary (σ) phases were

Fig. 9. (a) The power spectral density (PSD) profiles, (b), and (c) the multi-modal Gaussian histograms of topography results before and after immersion to thecorrosive medium, respectively. All of the analysis obtained from the AFM images in Fig. 8.

Fig. 10. Rainbow colour-map of AFM image. The correlation between surfaceconstituents and topography feature was shown which has been used for theextraction of data in Table 3.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

58

observed to be present in the 309/308 L stainless steel weld metal.2) The results obtained from the SVET study suggested that a number

of factors such as microstructure, composition, oxide film stabilityand time of exposure of the weldment to a corrosive environmentcan accelerate or decelerate the corrosion behavior of the dissimilarmetal weld. These factors play various roles in the galvanic inter-actions of the A508 base metal (anode) and the 309/308 L stainlesssteel weld metal (cathode) couple. The CGR region adjacent to thefusion boundary in A508 was observed to be the most susceptibleregion to corrosion, due to its complex microstructural surfaceconstituents (which are tempered bainite, martensite and MnS in-clusions).

3) AFM results analyzed with Gaussian multimodal distribution andfast Fourier transformation techniques clearly revealed quantita-tively the corrosion attacks on the surface constituents of the DMW.The trend of corrosion attack on the surface constituents is in theorder of: 309/308 L SS < weld borderline < tempered bainite inA508< tempered martensite in A508<MnS inclusion in A508.

4) A unified conclusion is drawn based on SVET, AFM, and SEM-EDSanalysis revealing that MnS induced pitting corrosion occurs si-multaneously with galvanic corrosion in A508 – 309/308 L SSweldment, where the former is more precarious. These corrosionprocesses identified could lead to a gradual degradation of the RPVsteel containers, and on a long run total damage of the RPV steel in

service. The implication of this damage is very severe to the NPP andthe environment, as it may lead to shut down of the NPP and pos-sible exposure of radioactive fission product to the environment.Thus, this galvanic corrosion understanding is a key to mitigatingthe potential damage of a RPV, especially at the bottom of the vesselwere parts of the SS cladding wears off due to mechanical damage.

Data availability

The raw/processed data required to reproduce these findings cannotbe shared at this time due to technical or time limitations.

Acknowledgements

This work is jointly supported by the National Key Research andDevelopment Program of China (2016YFE0105200, 2017YFB0702100),National Natural Science Foundation of China (No.51771211), KeyResearch Program of Frontier Sciences, CAS (QYZDY-SSW-JSC012) andthe Key Program of the Chinese Academy of Sciences (ZDRW-CN-2017-1).

References

[1] H.T. Wang, G.Z. Wang, F.Z. Xuan, C.J. Liu, S.T. Tu, Local mechanical properties of adissimilar metal welded joint in nuclear power systems, Mater. Sci. Eng.: A 568(2013) 108–117.

[2] M.-C. Kim, S.-G. Park, K.-H. Lee, B.-S. Lee, Comparison of fracture properties inSA508 Gr.3 and Gr.4N high strength low alloy steels for advanced pressure vesselmaterials, Int. J. Press. Vessels Pip. 131 (2015) 60–66.

[3] L. Li, C.F. Dong, K. Xiao, J.Z. Yao, X.G. Li, Effect of pH on pitting corrosion ofstainless steel welds in alkaline salt water, Constr. Build. Mater. 68 (2014) 709–715.

[4] E. Rahimi, A. Rafsanjani-Abbasi, A. Imani, S. Hosseinpour, A. Davoodi, Correlationof surface volta potential with galvanic corrosion initiation sites in solid-statewelded Ti-Cu bimetal using AFM-SKPFM, Corros. Sci. 140 (2018) 30–39.

[5] E. Rahimi, A. Rafsanjani-Abbasi, A. Imani, S. Hosseinpour, A. Davoodi, Insights intogalvanic corrosion behavior of Ti-Cu dissimilar joint: effect of microstructure andvolta potential, Mater. 11 (2018) 1820.

[6] M. Asemabadi, M. Sedighi, M. Honarpisheh, Investigation of cold rolling influenceon the mechanical properties of explosive-welded Al/Cu bimetal, Mater. Sci. Eng.: A558 (2012) 144–149.

[7] Z.R. Chen, Y.H. Lu, X.F. Ding, T. Shoji, Microstructural and hardness investigationson a dissimilar metal weld between low alloy steel and Alloy 82 weld metal, Mater.Charact. 121 (2016) 166–174.

[8] K.J. Choi, S.C. Yoo, T. Kim, C.B. Bahn, J.H. Kim, Effects of aging temperature onmicrostructural evolution at dissimilar metal weld interfaces, J. Nucl. Mater. 462(2015) 54–63.

[9] W. Wang, Y. Lu, X. Ding, T. Shoji, Microstructures and microhardness at fusionboundary of 316 stainless steel/Inconel 182 dissimilar welding, Mater. Charact. 107(2015) 255–261.

[10] Z. Wang, J. Xu, T. Shoji, Y. Takeda, H. Yuya, M. Ooyama, Microstructure and pittingbehavior of the dissimilar metal weld of 309L cladding and low alloy steel A533B, J.Nucl. Mater. 508 (2018) 1–11.

[11] N. Wint, J. Leung, J.H. Sullivan, D.J. Penney, Y. Gao, The galvanic corrosion ofwelded ultra-high strength steels used for automotive applications, Corros. Sci. 136(2018) 366–373.

[12] D. Liu, H. Yun Zhao, D. Yan Tang, Q. Ma, Corrosion behavior study of 22MnB5 steeland its weld using electrochemical method, Appl. Mech. Mater. 496 (2014)340–343.

[13] K.M. Deen, R. Ahmad, I.H. Khan, Z. Farahat, Microstructural study and electro-chemical behavior of low alloy steel weldment, Mater. Des. 31 (2010) 3051–3055.

[14] C. Garcia, F. Martin, P. de Tiedra, Y. Blanco, M. Lopez, Pitting corrosion of weldedjoints of austenitic stainless steels studied by using an electrochemical minicell,Corros. Sci. 50 (2008) 1184–1194.

Table 3Gaussian distribution parameters extracted from the topography histograms in Fig. 9b and c.

Region label Constituents Mean value of topography (μ ,nm) Standard deviation of topography (σ ,nm)

Before immersion LAS 508 80 5309/308 SS 55 10

After immersion MnS inclusions in LAS 508 350 35Martensite in LAS 508 500 70Tempered bainite in LAS 508 600 40Weld borderline 720 55309/308 L SS 860 50

Table 4Corrosion rate calculated from Table 3 and using Faraday law for major surfaceconstituents.

Surface constituents Calculated corrodeddepth (nm)

Calculated corrosion rate(μA cm−2)

MnS inclusions in LAS 508 800 ± 50 130 ± 10Martensite in LAS 508 420 ± 40 40 ± 5Tempered bainite in LAS

508270 ± 35 28 ± 3

Fig. 11. SEM image of AFM test sample after immersion in a corrosive mediafor 6 h, showing corrosion attack on the DMW; the elemental composition of thepit marked with a cross is provided in the inserted table at the top corner.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

59

[15] S.J. Zinkle, G.S. Was, Materials challenges in nuclear energy, Acta Mater. 61 (2013)735–758.

[16] Light Water Reactor Sustainability Program, Integrated Program Plan, INL/EXT -11-23452, Office of Nuclear Energy, US Department of Energy, 2012.

[17] G.R. Odette, R.K. Nanstad, Predictive reactor pressure vessel steel irradiation em-brittlement models: issues and opportunities, JOM 61 (2009) 17–23.

[18] http://www.chuden.co.jp/corporate/publicity/pub_release/press/3258691_21432.html.

[19] J. Soltis, Passivity breakdown, pit initiation and propagation of pits in metallicmaterials – review, Corros. Sci. 90 (2015) 5–22.

[20] T. Kondo, H. Nakajima, R. Nagasaki, Metallographic investigation on the claddingfailure in the pressure vessel of a BWR, Nucl. Eng. Des. 16 (1971) 205–222.

[21] S.-W. Kim, D.-J. Kim, H.-P. Kim, Evaluation of galvanic corrosion behavior of Sa-508 low alloy steel and type 309l stainless steel cladding of reactor pressure vesselunder simulated primary water environment, Nucl. Eng. Technol. 44 (2012)773–780.

[22] S. Wang, H. Ming, J. Ding, Z. Zhang, J. Wang, E.-H. Han, A. Atrens, Effect of H3BO3

on corrosion in 0.01 M NaCl solution of the interface between low alloy steel A508and alloy 52 M, Corros. Sci. 102 (2016) 469–483.

[23] S. Wang, J. Ding, H. Ming, Z. Zhang, J. Wang, Characterization of low alloy ferriticsteel–Ni base alloy dissimilar metal weld interface by SPM techniques, SEM/EDS,TEM/EDS and SVET, Mater. Charact. 100 (2015) 50–60.

[24] S. Weng, Y. Huang, F. Xuan, F. Yang, Pit evolution around the fusion line of aNiCrMoV steel welded joint caused by galvanic and stress-assisted coupling corro-sion, RSC Adv. 8 (2018) 3399–3409.

[25] A.I. Karayan, H. Castaneda, Weld decay failure of a UNS S31603 stainless steelstorage tank, Eng. Fail. Anal. 44 (2014) 351–362.

[26] L.W. Wang, Z.Y. Liu, Z.Y. Cui, C.W. Du, X.H. Wang, X.G. Li, In situ corrosioncharacterization of simulated weld heat affected zone on API X80 pipeline steel,Corros. Sci. 85 (2014) 401–410.

[27] F. Mohammadi, F.F. Eliyan, A. Alfantazi, Corrosion of simulated weld HAZ of API X-80 pipeline steel, Corros. Sci. 63 (2012) 323–333.

[28] Y. Zhao, J. Dong, Y. Ma, L. Zhao, X. Pei, Mechanical and galvano-chemistry prop-erty variation within dissimilar metal weld between 1Cr18Ni9Ti and 1Cr13 stain-less steel, J. Mater. Sci. Technol. 26 (2010) 477–480.

[29] Z.F. Yin, M.L. Yan, Z.Q. Bai, W.Z. Zhao, W.J. Zhou, Galvanic corrosion associatedwith SM 80SS steel and Ni-based alloy G3 couples in NaCl solution, Electrochim.Acta 53 (2008) 6285–6292.

[30] Q. Qiao, G. Cheng, W. Wu, Y. Li, H. Huang, Z. Wei, Failure analysis of corrosion atan inhomogeneous welded joint in a natural gas gathering pipeline considering thecombined action of multiple factors, Eng. Fail. Anal. 64 (2016) 126–143.

[31] Z. Wang, B. Liu, Y. Yang, N. Han, G. Zeng, A. Ren, T. Zhang, Experimental andnumerical studies on corrosion failure of a three-limb pipe in natural gas field, Eng.Fail. Anal. 62 (2016) 21–38.

[32] A. Stenta, S. Basco, A. Smith, C.B. Clemons, D. Golovaty, K.L. Kreider, J. Wilder,G.W. Young, R.S. Lillard, One-dimensional approach to modeling damage evolutionin galvanic corrosion, Corros. Sci. 88 (2014) 36–48.

[33] C.A. Huang, T.H. Wang, W.C. Han, C.H. Lee, A study of the galvanic corrosionbehavior of Inconel 718 after electron beam welding, Mater. Chem. Phys. 104(2007) 293–300.

[34] H. Ming, Z. Zhang, J. Wang, E.-H. Han, W. Ke, Microstructural characterization ofan SA508–309L/308L–316L domestic dissimilar metal welded safe-end joint,Mater. Charact. 97 (2014) 101–115.

[35] Q. Wang, M. Zhang, W. Liu, X. Wei, J. Xu, J. Chen, H. Lu, C. Yu, Study of type-IIboundary behavior during SA508-3/EQ309L overlay weld interfacial failure pro-cess, J. Mater. Process. Technol. 247 (2017) 64–72.

[36] Q. Xiong, H. Li, Z. Lu, J. Chen, Q. Xiao, J. Ma, X. Ru, Characterization of micro-structure of A508III/309L/308L weld and oxide films formed in deaerated high-temperature water, J. Nucl. Mater. 498 (2018) 227–240.

[37] C.-C. Hsieh, W. Wu, Overview of intermetallic sigma (σ) phase precipitation instainless steels, ISRN Metall. 2012 (2012) 1–16.

[38] M.H. Lewis, Precipitation of (Fe, Cr) sigma phase from austenite, Acta Metall. 14(1966) 1421–1428.

[39] N. Lopez, M. Cid, M. Puiggali, Influence of sigma phase on mechanical propertiesand corrosion resistance of duplex stainless steels, Corros. Sci. 41 (1999)1615–1631.

[40] A. Davoodi, Z. Esfahani, M. Sarvghad, Microstructure and corrosion characteriza-tion of the interfacial region in dissimilar friction stir welded AA5083 to AA7023,

Corros. Sci. 107 (2016) 133–144.[41] H. Ming, J. Wang, E.-H. Han, Comparative study of microstructure and properties of

low-alloy-steel/nickel-based-alloy interfaces in dissimilar metal weld joints pre-pared by different GTAW methods, Mater. Charact. 139 (2018) 186–196.

[42] O.V. Karavai, A.C. Bastos, M.L. Zheludkevich, M.G. Taryba, S.V. Lamaka,M.G.S. Ferreira, Localized electrochemical study of corrosion inhibition in micro-defects on coated AZ31 magnesium alloy, Electrochim. Acta 55 (2010) 5401–5406.

[43] J.A. Moreto, C.E.B. Marino, W.W. Bose Filho, L.A. Rocha, J.C.S. Fernandes, SVET,SKP and EIS study of the corrosion behaviour of high strength Al and Al–Li alloysused in aircraft fabrication, Corros. Sci. 84 (2014) 30–41.

[44] T.G. Thompson, W.R. Johnston, H.E. Wirth, The sulfate chlorinity ratio in oceanwater, Conseil Perm. Internat. p. l’Explor, de la Mer, Jour.du Conseil 6 (1931)246–251.

[45] A.L. Shaeffler, Constitution diagram for stainless steel weld metal, Met. Prog. 56(1949) 680-680B.

[46] F. Mas, M. Guilhem, L. Pierre, B. Yves, T. Catherine, F. Roch, T. Patrick, A. Simar,Heterogeneities in local plastic flow behavior in a dissimilar weld between low-alloy steel and stainless steel, Mater. Sci. Eng.: A 667 (2016) 156–170.

[47] K.-H. Lee, S.-g. Park, M.-C. Kim, B.-S. Lee, D.-M. Wee, Characterization of transitionbehavior in SA508 Gr.4N Ni–Cr–Mo low alloy steels with microstructural alterationby Ni and Cr contents, Mater. Sci. Eng.: A 529 (2011) 156–163.

[48] G. Green, R. Higginson, S. Hogg, S. Spindler, C. Hamm, Evolution of sigma phase in321 grade austenitic stainless steel parent and weld metal with duplex micro-structure, Mater. Sci. Technol. 30 (2013) 1392–1398.

[49] S. Kou, Welding Metallurgy, second edition, John wiley & sons, Inc., 2019 ISBN: 0-471- 43491-4 127.

[50] C.-C. Hsieh, W. Wu, Discussing the precipitation behavior of σ phase using diffusionequation and thermodynamic simulation in dissimilar stainless steels, J. AlloysCompd. 506 (2010) 820–825.

[51] C. Aya, M. Izumi, S. Yu, H. Nobuyoshi, Pit initiation mechanism at MnS inclusion instainless steel: synergistic effect of elemental sulfur and chloride ions, J.Electrochem. Soc. 160 (2013) C511–C520.

[52] L.-B. Niu, K. Okano, S. Izumi, K. Shiokawa, M. Yamashita, Y. Sakai, Effect ofchloride and sulfate ions on crevice corrosion behavior of low-pressure steam tur-bine materials, Corros. Sci. 132 (2018) 284–292.

[53] A.G. Marques, J. Izquierdo, R.M. Souto, A.M. Simões, SECM imaging of the cut edgecorrosion of galvanized steel as a function of pH, Electrochim. Acta 153 (2015)238–245.

[54] B. Stellwag, The mechanism of oxide film formation on austenitic stainless steels inhigh temperature water, Corros. Sci. 40 (1998) 337–370.

[55] U. Donatus, G.E. Thompson, H. Liu, X. Zhou, Z. Liu, Understanding the galvanicinteractions between AA2024T3 and mild steel using the scanning vibrating elec-trode technique, Mater. Chem. Phys. 161 (2015) 228–236.

[56] J. Izquierdo, L. Martín-Ruíz, B.M. Fernández-Pérez, L. Fernández-Mérida,J.J. Santana, R.M. Souto, Imaging local surface reactivity on stainless steels 304 and316 in acid chloride solution using scanning electrochemical microscopy and thescanning vibrating electrode technique, Electrochim. Acta 134 (2014) 167–175.

[57] A.M. Simões, D. Battocchi, D.E. Tallman, G.P. Bierwagen, SVET and SECM imagingof cathodic protection of aluminium by a Mg-rich coating, Corros. Sci. 49 (2007)3838–3849.

[58] N. Sathirachinda, R. Pettersson, J. Pan, Depletion effects at phase boundaries in2205 duplex stainless steel characterized with SKPFM and TEM/EDS, Corros. Sci. 51(2009) 1850–1860.

[59] N. Sathirachinda, R. Pettersson, S. Wessman, J. Pan, Study of nobility of chromiumnitrides in isothermally aged duplex stainless steels by using SKPFM and SEM/EDS,Corros. Sci. 52 (2010) 179–186.

[60] X. Wu, Y. Katada, I.S. Kim, S.G. Lee, Hydrogen-involved tensile and cyclic de-formation behavior of low-alloy pressure vessel steel, Metall. Mater. Trans. A 35(2004) 1477–1486.

[61] Z. Esfahani, E. Rahimi, M. Sarvghad, A. Rafsanjani-Abbasi, A. Davoodi, Correlationbetween the histogram and power spectral density analysis of AFM and SKPFMimages in an AA7023/AA5083 FSW joint, J. Alloys Compd. 744 (2018) 174–181.

[62] E. McCafferty, Introduction to Corrosion Science, Springer science & businessmedia, 2010, p. 23 ISBN: 978-1-4419-0454-6.

[63] A.C. Bastos, M.C. Quevedo, O.V. Kuravai, M.G. Ferreira, Review-On the applicationof the scanning vibrating electrode techniques (SVET) to corrosion research, J.Electrochem. Soc. 164 (2017) C973–C990.

B.O. Okonkwo, et al. Corrosion Science 154 (2019) 49–60

60