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Corrosion behavior of stainless steel weldments inphysiological solutionsTo cite this article: A Farooq et al 2018 Mater. Res. Express 5 015401

 

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Mater. Res. Express 5 (2018) 015401 https://doi.org/10.1088/2053-1591/aa9f98

PAPER

Corrosion behavior of stainless steel weldments in physiologicalsolutions

AFarooq1 ,MAzam2 andKMDeen1,3

1 CorrosionControl ResearchCell, Department ofMetallurgy andMaterials Engineering, CEET,University of the Punjab, 54590, Lahore,Pakistan

2 Institute of Chemical Engineering&Technology, University of the Punjab, 54590, Lahore, Pakistan3 Department ofMaterials Engineering, University of British Columbia, Stores Road, V6T 1Z4Vancouver, BC, Canada

E-mail: kmdeen.ceet@pu.edu.pk

Keywords: stainless steel, welding, corrosion

AbstractIn this study corrosion behavior of TIGwelded 316L stainless steel plates in simulated biologicalsolutions is investigated. Themechanical testing results showed slight decrease in ductility afterwelding and the fracture surface representedmixed cleavage and inclusions containing dimplestructure. The heat affected andweld zone (WZ) demonstrated higher corrosion potential andrelatively large pitting tendency than basemetal (BM) in bothHank’s andRinger’s solution. Theformation of delta (δ) ferrite in the heat affected andWZdecreased the corrosion resistance asconfirmed frompotentiodynamic Tafel scans. The decrease in pitting resistance and lower protectiontendency of theWZ compared to BMand heat affected zonewas also quantified from the cyclicpolarization trends.

1. Introduction

Stainless steel, Co–Cr,Ni–Cr and titaniumbased alloys arewidely used in orthopedics as implantsmaterials dueto their goodmechanical properties, chemical stability in the biological environments and low cost [1–3]. Thebiological environment is composed of water, dissolved oxygen, salts, organic compounds and complex aminoacid structure. These ingredients could affect the functionality of bio-implants adversely if the surface of theimplant has structural and design defects. Therefore fabrication, effective post treatment procedures, andexcellent surgical practice could enhance the longevity of the implantmaterials. Stainless Steel AISI 316L is oneof themost commonly usedmaterial in the orthopedics due to its economical fabrication, goodmechanicalproperties [3, 4] and inherent tendency to form compact and self-healing passive film in the oxidizingenvironments [1]. The functionality of the surrounding tissuesmay also be adversely affected due to the releaseof nickel and chromium ions exceeding tolerance limits. These ions could initiate toxic, allergic or carcinogenicreactions in the surrounding tissues andmay sometimes be life threatening [5].

Various components of themetallic implants are joined throughweldingwhich could introducemetallurgical defects if the quality procedures are not employed. Although austenitic stainless steels have goodweldability but thermal spikes duringweldingmay altermicrostructure which could influence themechanicaland electrochemical properties. Also the inherent compositional andmicrostructural variation in theweld andthermally affected zone could also change the electrochemical behavior of the alloys [6, 7]. Prosthetic devices e. g.mono-bloc hip stems aremade of stainless steel and arewelded not because offixation requirements but due totechnical and commercial needs [8, 9]. These devices are designed to stay in the body for long periods and anymanufacturing defect could lead to premature failure. This scenario is always painful and risky for a patient andrequires definite repetitive post-surgical procedures [10]. The aggressive nature of body environment andformation of electrochemical cells at the tissue implant interface could accelerate the localized corrosionreactions [11].

RECEIVED

30August 2017

REVISED

21November 2017

ACCEPTED FOR PUBLICATION

6December 2017

PUBLISHED

3 January 2018

© 2018 IOPPublishing Ltd

The objective of this studywas to determine the effect of welding on themicrostructural variation andmechanical properties of the AISI 316L. Also the electrochemical behavior of basemetal (BM), heat affected andweld zone (WZ)was examined in the simulated physiological environment i.e. Ringer’s andHank’s solutions atambient temperature.

2.Materials andmethods

Twoplates of 316L austenitic stainless steels (152.4 mm×152.4 mm×4.5 mm)were degreased in a soapsolution followed by cleaning in acetone.Weld joint was prepared by beveling the edge of one plate to single ‘V’shapewith an included angle of 45°. These plates were thenwelded byTungsten Inert Gas (TIG) technique using99.95%pureArgon to prevent oxidation of weld pool. The 3 mmdiameter wire (AWS-A5.9ER316)was used as afiller tomake a Butt joint. The composition of BM,filler wire andwelding parameters are given in tables 1 and 2respectively.

The three samples for tensile testingwerewire cut as per standard (ASTME8-03) and 1 cm2 samples formetallographywere sectioned from each zone i.e. BM, thermally/heat affected zone (HAZ), andWZ. Theschematic diagramof thewelded plate is shown in figure 1which describes the location of each sample cut fromthe BM,HAZ andWZ formicrostructural,mechanical and electrochemical analyses. The fracture surface aftertensile testingwas examined through scanning electronmicroscope (SEM). Formicrostructural study onesurface from each zonewas ground using silicon carbide papers up to 1200 grit size. Final polishingwas donewith 1 μmdiamond paste in order to produce scratch-freemirror-finished surface. The samples were degreasedwith acetone andwashedwith alkaline solution followed by rinsing in the distilledwater. These samples wereetched in a solution containing 3 parts hydrochloric acid (HCl), 2 parts glycerol and 1 part of nitric acid (HNO3)to reveal themicrostructural details through opticalmicroscope.

Themicrostructure and variation in the chemical composition of theWZwas examined through SEMandenergy dispersive x-ray (EDX) analysis, respectively. X-ray diffraction (XRD)patterns of BMandWZwere alsoobtained by usingCuKα1 (λ=1.5405 A°) as radiation source. The hardness of thewelded joint wasdetermined across theweld line according toASTME18-03 standard to elaborate themicrostructural features.

For electrochemical analysis; working electrodes (1.44 cm2) fromBM,HAZ andWZwere ground to 1000grit followed by soldering a glass encased copper wire at the back of each specimen. Thesewere then coldmounted in a polyester resin by exposing one surface to the physiological solutions at room temperature. Thechemical composition of both physiological solutions i.e. Ringer’s andHank’s is given in table 3.

Potentiodynamic Tafel scan and cyclic polarization (CP) testmethodswere used to investigate the corrosionkinetics, passive behavior and pitting tendency of each zone in these solutions. SaturatedCalomel was used asreference and compacted graphite rodwas used as counter electrode in a three electrode cell system connectedwithGamry Potentiostat (PC14-750). All the experiments were repeated three times and the average value ofeach parameter is given in this study.

3. Results and discussion

3.1.Microstructural study ofweldmentThemicrostructural features of BM,HAZ andWZdepend on the thermal excursion duringwelding process.The BMrepresented the relatively finer grains of austenite (γ-phase), represented as ‘A’ infigure 2(a) comparedto austenite in theHAZ (figure 2(b)) near fusion line. This behavior was attributed to the recrystallization ofgrain structure during thermal excursion. The pseudo-equilibrium conditions duringwelding, the heating and

Table 1.Chemical composition of the basemetal (316L) stainless steel andweldfillermetal (inwt%).

Material C Mn Cr Ni Mo P S Fe

Basemetal 0.025 1.20 16.74 9.37 2.28 0.065 0.062 Bal.

Fillermetal 0.022 1.35 17.00 11.52 2.10 0.031 0.014 Bal.

Table 2.The parameter used in TIGwelding process to join two plates of 316L stainless steel.

Material

Welding

method Fillermaterial

Diameter offillermateri-

als (mm)Welding cur-

rent (A)Welding

Groove

Welding

inert gas

316L TIG AWSA5.9ER316 3 150 ‘V’ Argon

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

cooling cycle also resulted in the formation of delta (δ) ferrite phase designated as ‘F’. In BM,HAZ andWZ the δ-ferrite phasewas elongatedwithin the austenitematrix (γ) depicting the direction of heatflow.However, theconcentration of δ-ferrite was higher inWZ thanBMandHAZ as shown infigure 2(c). The rapid cooling ofWZduringwelding cycle restricted the complete peritectic transformation of δ-ferrite into austenite (γ) [12]. Thesolidification ofWZ followed the [L→L+δ→L+δ+γ→γ+δ] sequence representing the nucleationof δ-ferrite in the liquid (L)weldmetal followed by the transformation of L-phase into the austenite (γ). Theγ-phase formation in theWZ followed the nucleation of δ-ferrite during rapid solidification of themoltenmetal intheweld pool. To further elaborate the existence of δ-ferrite in theγ-phase, the high resolution image ofWZ isshown infigure 3(a). Due to non-equilibrium cooling conditions in theWZ and heterogeneous solid-state phasetransformation, the complete conversion of δ-ferrite intoγ-phasewas limited. By comparing the variation ofalloying elements i.e. Cr, Ni,Mn andMo concentrations in the δ-ferrite andγ-phase, it is possible to predictreason of δ-ferrite formation in theWZ. It was determined that the concentration of ferrite stabilizing elementsi.e. Cr andMowas low in the δ-ferrite compared toγ-phase as given in the table offigure 3. This could possiblybe relatedwith the hindered diffusion of these elements from themoltenmetal in theweld pool to the δ-ferritenuclei during transient welding conditions.

Figure 1. Schematic diagramof thewelded plate showing the location of samples cut formechanical,microstructural andelectrochemical testing.

Table 3.Compositions of the physiological solutions (in g l−1).

Components

Hank’s solu-

tion pH=7.4Ringer’s solu-

tion pH=7.0

NaCl 8.00 8.5

KCl 0.40 0.2

CaCl2 0.14 0.2

NaHCO3 0.35 0.2

MgSO4.7H2O 0.06 —

MgCl2.6H2O 0.10 —

KH2PO4 0.06 —

Glucose 1.00 —

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

Figure 2.Opticalmicrograph of AISI 316Lweldment showing the distribution of δ-ferrite in the (a)BM (b)HAZand (c)WZ.

Figure 3. SEMandEDX analysis of the austeniticmatrix and ferrite phases (δ-ferrite) in theWZ.

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

Also, the δ-ferrite in theWZhad acute angle to the ferrite in theHAZ andBMacross the fusion line. Thiscould be simulatedwith the swirlingmotion ofmoltenweld pool due to arc striking and its forwardmotionalong theweld line.

The formation of δ-ferrite could deleteriously affect themechanical and electrochemical properties ofstainless steel. The higher concentration of δ-ferrite could disrupt the formation of uniformpassive film over theausteniticmatrix (A) resulting in the initiation of localized corrosion reactions [13, 14]. In other way theexistence of δ-ferrite in theWZwould be advantageous tominimize hot cracking tendency duringsolidification [15].

The presence of δ-ferrite in the austenitematrix (inWZ)was further confirmed fromXRDpatterns. Figure 4represents the diffraction pattern of BMandWZ. The broad peaks originated at 47.4° and 64.8° in the diffractionpattern ofWZ corresponded to the (110) and (200) planes of δ-ferrite. On the other hand the peaks affiliatedwith the austenite (γ)were centered at 2θ=43.2°, 50.5° and 74.8°which represented the (111), (200) and (220)crystallographic planes, respectively of thematrix phase. The evidence of otherNi andMn rich phase (i.e.Ni2Mn3Si) peaks in the XRDpattern ofWZ also supported the relatively high concentration of these elements inδ-ferrite as confirmed by the EDX analysis (figure 3).

3.2.Mechanical testing and fractographyThe hardness profile of theweldment fromBM toWZ is shown infigure 5. It is represented that the averageRockwell Hardness of BMwas 78.5±0.3HRBwhich gradually increased to 83.3±0.5HRB at the fusion line.In the center ofWZ the hardness wasmuch higher (85.3±0.5HRB) than in theHAZ andBM.The increase inhardness was relatedwith the high concentration of δ-ferrite in the austenitematrix [8]. As δ-ferrite (BCCphase)was non-coherent with the γ-matrix (FCC) andmay increase the hardness inHAZ andWZ compared to BM.

The tensile testing of welded specimenswas carried out in triplicate by using Instron 1120 universal testingmachine. The decrease in the ductility of welded sample than 316L (BMonly, withoutwelding)was attributed tonon-homogeneity in themicrostructure (existence of δ-ferrite)which resulted in the fracture of tensile samplefromWZ. TheweldedAISI 316L samples showed relatively higher hardness, yield strength (YS) and ultimatetensile strength thanBMbut these values were comparable to the fillermetal (AWSA5.22 E316) properties astabulated in table 4.

The 30% elongation of welded 316L samples before fracturewas lower than un-weldedAISI 316L (40%)steel and for comparison, it is recommended to see [16]. The fracture surface showed typical ductile dimple likestructuremostly formed by the coalescence ofmicro voids originated at the inclusions as evident infigure 6.

The shallow shear tongue of dimples corresponded to the increase in hardness after weldingwhichwas alsoin support with lowpercentage elongation. It was noticed that some intergranularmicro-cracks were presentaround the carbide precipitates within theWZ. The existence of transgranular crack along the crystallographicplanesmay produce cleavage facets alongwith dimplemorphology as evident infigure 6(b). Themixedmode offracture behavior through theWZwas attributed to the formation of non-coherent δ-ferrite in the austeniticmatrix. The presence of second phase within themicrostructure also validated the coalescence ofmicro-voids

Figure 4.X-ray diffraction patterns of the basemetal (BM) andweld zone (WZ).

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

leading tofinal fracture due to segregation of δ-ferrite along the grain boundaries during rapid cooling ofmoltenmetal in theweld pool.

The composite (δ+γ)microstructure ofWZmay also result in the decrease in ductility than the AISI 316Lun-weldedmaterial. It was deduced that non-equilibrium thermalfluctuations and cooling affected themechanical properties of theweldedAISI 316L. This behaviorwas relatedwith the presence of largeconcentration of δ-ferrite and other phases in theWZ.

3.3. Electrochemical investigation ofweldments3.3.1. Open circuit potential (OCP)TheOCPof BM,HAZ andWZ in bothRinger’s andHank’s solutionswas determinedwith respect to Saturatedcalomel (SCE) reference electrode (+250 mVversus SHE) as given in table 5. It was observed thatOCPof BM,HAZ andWZ inRinger’s solutionwasmore negative than inHank’s solution. This was attributed to the highchloride contents in the Ringer’s solution, whichmay possibly decrease theOCP compared toHank’s solution.At pH (7.4±0.1) the stable species i.e. -H PO2 4

1 and -HPO42 in the hanks solution have the ability to

preferentially interact with the surface of stainless steel. These species could promote the stability of passive filmover the surface andmay limit the localized corrosion tendency of 316L stainless steel [1, 6]. The relativelynegativeOCP value ofHAZ inRinger’s solutionwas considered due to the recrystallization ofγ-phase andnucleation of δ-ferrite along the grain boundaries. These sites could be the active sites for electrochemical attackat the surface. BothHAZ andWZwere found to bemore anodic (negativeOCP) compared to the BM inHank’ssolutionwhereas, thermodynamically theHAZwasmore active than BMandWZ in the Ringer’s solution.

3.3.2. Potentiodynamic polarization scansTypical Potentiodynamic Tafel Scans for BM,HAZ andWZwere obtained (figure 7) in both simulatedphysiological solutions. The polarization behavior was evaluated in the potential range of (−0.5 to 1.5 V)OCPwith a scan rate of 1 mV s−1. The kinetic parameters of BM,HAZ andWZwere calculated by Tafel extrapolationmethod using EchemAnalyst software as given in table 6. The higher corrosion current density (Icorr) ofHAZandWZ in ringer solution represented higher corrosion rate compared to BMaccording toButler–Volmerrelation [17].

Figure 5.Hardness profile of AISI 316weldment 5.3 cmaway fromweld center line.

Table 4.Mechanical properties of the 316Lweldment.

Samples Yield strength (MPa) Tensile strength (MPa) %Elongation

Welded sample 377.842±5.0 525.095±4.0 30±1.0AWSA5.22 E316 400 560 35

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

b h b h= - - -⎜ ⎟⎧⎨⎩

⎛⎝

⎞⎠

⎡⎣⎢

⎤⎦⎥

⎫⎬⎭( ) ( )i inF

RT

nF

RTexp 1exp , 1onet a a c c

Figure 6. SEM images of fractured surface after tensile testing (a) ductile dimple structure (b)mixedmode fracture(dimple+ cleavage).

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

Table 5.Open circuit potential (OCP) in both physiological solutions.

Samples BM HAZ WZ

Hank’s Solution

OCP (mV) −139.3±5 −292.09±8 −220.09±7Ringer’s Solution

OCP (mV) −234.75±7 −368.33±8 −239.25±10

Figure 7.Potentiodynamic polarization scans (a)Ringer’s (b)Hank’s solutions.

Table 6.Potentiodynamic Tafel scans extrapolation parameters.

Samples βa (mV/decade) βc (mV/decade) Icorr (μA cm−2) Ecorr (mV) Corrosion rate (mpy)

Ringer’s Solution

BM 298.5 93.50 0.241±0.01 −228.0±10 0.104±0.04HAZ 1876 217.6 2.53±0.02 −398.0±6 1.102±0.05WZ 1745 149.3 2.15±0.05 −275.0±8 0.934±0.04

Hank’s Solution

BM 219.4 72.70 0.0715±0.005 −187.0±6 0.031±0.006HAZ 253.5 82.6 0.0842±0.004 −292.0±9 0.036±0.002WZ 349.9 120.2 0.450±0.008 −277.0±12 0.195±0.02

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

where ‘ba’ and ‘bc’ are the anodic and cathodic Tafel constants. ‘ha’ and ‘hc’ represent the anodic and cathodicpolarization of themetal surfaces during Tafel extrapolation.

The presence of δ-ferrite in the austeniticmatrix could disrupt the uniformity of passive film. Thepreferential dissolution of δ-ferrite caused by the formation of local galvanic cells within the austeniticmatrixcould enhance the pitting tendency of surface in the chloride containing physiological solution. The higher Icorrandmore negative (active) corrosion potentialEcorr ofHAZ thanBM in the Ringer’s solutionwas also attributedto the formationmicrostructural inhomogeneity duringwelding process. TheWZ zone also presented highcorrosion rate than BM in the same solution andwas relatedwith the relatively high concentration of dispersedδ-ferrite in the austenitic phase. However, under steady state condition (at or nearOCP), the relative smalldifference in the Ecorr ofWZ (−275.0 mV)SCE andBM (−228.9 mV)SCEmay be less deleterious in terms of localgalvanic cell formation.

InHank’s solution the corrosion rate ofWZwas relatively high and followed the trend (corrosion rate ofBM<HAZ<WZ). In comparison, the corrosion rate of AISI 316Lweldment (of all BM,HAZ andWZ) in theRinger’s solutionwas higher than inHank’s solution. The lower corrosion rate inHank’s solutionmay be relatedwith the adsorption of stable species ( -H PO2 4

1 -HPO42 ) at the local anodic sites developed at the surface under

applied conditions. These adsorbed anionsmay restrict the de-polarization of surface and hencemay decreasethe ingress of other aggressive ions towards the surface. In relationwith adsorption, the electrochemical stabilityof the passivefilmwas enhanced and therefore decreased the corrosion tendency of weldment inHank’ssolution. But the relatively high concentration of δ-ferrite in theWZwithin and along the grain boundariescould disrupt the uniformity of the passive film,which could be the possible reason of high corrosion rate inboth solutions. The reported results about the electrochemical behavior of are in support and comparable toother electrochemical investigations on literature austenitic steel welded joints [12, 18–20]. To evaluate thepassivation tendency of different zones in the Ringer’s solution, the each samplewas polarized from itsOCP to+1.5 V (versusOCP)with a scan rate of 5 mV s−1. The passive current density (ip) ofHAZwas 35.93 μA cm−2

compared to 13.80 and 15.05 μA cm−2 for BMandWZ, respectively. Similarly in theHank’s solution, the ip forHAZ (11.03 μA cm−2)was slightly higher than BM (9.59 μA cm−2) butWZpresented no limiting current value(representing passive behavior). The continuous increase in the current density during anodic polarizationwasassociatedwith formation of relatively unstable passive filmdue to presence of large concentration δ-ferrite inthe γmatrix.

3.3.3. CP characterizationDuringCP the surface of all samples was scanned between−0.5 mV/OCP to an apex potential (1.00 V versusOCP). The pitting resistance and protection tendencywasmeasured by reversing the potential from the apexpotential at a scan rate of 2.5 mV s−1. The abrupt increase in the current during forward (anodic) scan belongs tothe transpassive region corresponding to the breakdownof passive film, designated as pitting potential, Epit. Theintersection of reverse polarization curvewith the anodic polarization line resulted in the positive hysteresisloop, represented the pitting tendency and is designated as protection potentials, Eprot. In otherwords, thepittingmay happen aboveEprot and can be used as a diagnostic to estimate complete protection of surface fromlocalized attack under applied conditions. There is always a possibility for the formation ofmetastable pits abovethis potential which are difficult to analyze by other physical characterizationmeans. However, the potentiallarger thanEpit represents the formation of stable pits at the surface.

TheCP scans of all the samples are shown infigure 8 and the values of corrosion potential, (Ecorr) pittingpotential (Epit) and protection potential (Eprot) are given in table 7.

The stability of passive film in theRinger’s andHank’s solution depends on themicrostructure andelectrochemical species present in the electrolytes. Pit propagation kinetics wasmeasured by analyzing thehysteresis loop in theCP trends [13, 21, 22].

The anodic polarization depicted the rapid increases in current, which reflected the pit initiation and growthdue to breakdownof passive film from the localized sites. In Ringer’s solution, the small current fluctuationwithin the passive region ofHAZ andWZwas observedwhich confirmed to the presence of localized active siteswithin themicrostructure, corresponding to themetastable pitting tendency. These localized corrosionreactions at the surface were attributed to the presence δ-ferrite in the austenitic (γ)matrix.

The lower pitting resistance and protection tendency ofHAZ andWZ thanBM in both solutions arerepresented infigure 9. In comparison to the solution types, the presence of -H PO2 4

1 and -HPO42 inHank’s

solution behaved as buffer tomodify the passivefilm andmay adsorb at the local active sites to restrict theinitiation and formation of pits at the surface [23].

The higher pitting resistance of BM thanHAZ andWZdemonstrated the direct relationwith themicrostructural inhomogeneity produced duringweld thermal cycle. The decreasing trend in the pittingresistance fromBM toWZ indicated the active localized dissolution ofHAZ andWZ. The pitting resistance andprotection tendency of welded samples in Ringer’s solutionwas lower than inHank’s solutionwhich could be

9

Mater. Res. Express 5 (2018) 015401 A Farooq et al

associatedwith the high activity of Cl− ions species in the former electrolyte. Also, the -H PO2 41 and -HPO4

2

species present in theHank’s solution could preferentially adsorb at the active sites andmay possibly restrict thedissolution tendency. The variation in the δ-ferrite concentration and other phases in themicrostructure ofHAZ andWZ could deteriorate the uniformity of passive film. It has been evaluated that the heterogeneity in themicrostructure duringwelding could deleteriously affect the corrosion properties of AISI 316L andmayaccelerate localized dissolution ofWZ in the simulated physiological solution.

Figure 8.Cyclic polarization scans for AISI 316L (a) in Ringer’s (b)Hank’s solutions.

Table 7.Electrochemical parameters calculated fromCyclic Polarization curves, the average values of each parameter is given form thetriplicate test with a deviation of±5%.

Samples Ecorr (mV) Eprot (mV) Epit (mV) Pitting resistance=Epit−Ecorr (mV) Protection tendency=Eprot−Ecorr (mV)

Ringer’s Solution

BM −376.3 −53.76 69.89 446.19 322.54

HAZ −338.7 −104.8 48.39 387.09 233.9

WZ −306.5 −123.7 16.13 322.63 182.8

Hank’s Solution

BM −346.2 221.9 357.7 703.9 568.1

HAZ −461.5 −35.5 147.9 609.4 426.0

WZ −346.2 −65.09 97.63 248.57 281.11

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

4. Conclusion

1. This study concluded that welding thermal cycle greatly affected the microstructural, mechanical andelectrochemical characteristics of the AISI316Lweldment.

2. The formation of δ-ferrite and other phases in theWZ were confirmed from EDX and XRD analysis whichpossibly deteriorate themechanical properties of theweldment by decreasing%elongation and fracturingthrough theWZduring tensile test experiments.

3. The fracture surface morphology also revealed mixed mode (Dimple+Cleavage) type of failure due tomicrostructural features ofWZ.

4. The higher concentration of δ-ferrite and formation other phases in the WZ deleteriously affected theelectrochemical properties of theweldment.

5. The relatively higher corrosion rate, lower pitting resistance and protection tendency of WZ than BM wasconfirmed from the potentiodynamic andCP analyses.

6. The Ringer’s solution in contrast to Hank’s solution was more aggressive towards BM, HAZ and WZ. Thepresence of -H PO2 4

1 and -HPO42 species in theHank’s solutionmay preferentially adsorb at the active

sites andmay possibly hinder the pitting corrosion of AISI 316Lweldment.

ORCID iDs

AFarooq https://orcid.org/0000-0003-0663-4234

Figure 9.Electrochemical properties of weldments evaluated from the cyclic polarization trends (a) pitting protection (b) protectiontendency; note: the values reported here are the average valueswith standard deviation of±5%.

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Mater. Res. Express 5 (2018) 015401 A Farooq et al

KMDeen https://orcid.org/0000-0002-3619-2599

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