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Microstructure and corrosion behavior of multipass gas tungsten arc welded 304L stainless steel G.R. Mirshekari, E. Tavakoli, M. Atapour , B. Sadeghian Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran article info Article history: Received 11 June 2013 Accepted 24 October 2013 Available online 5 November 2013 Keywords: Gas tungsten arc welding Microstructure Corrosion resistance abstract The purpose of this study was to discuss the effect of single pass and multipass (double and triple pass) gas tungsten arc welding (GTAW) on microstructure, hardness and corrosion behavior of 304L stainless steel. In this investigation, 308 stainless steel filler metal was used. Microstructures and hardness of the weldments were investigated using optical microscopy, scanning electron microscopy (SEM), energy dis- persive spectroscopy (EDS), X-ray diffraction analysis (XRD) and Vickers microhardness (HV 0.5 ). A ferrite- scope was also used in the non-destructive evaluation to observe the ferrite content on the weldments. Corrosion behavior of weldments in 1 M H 2 SO 4 solution at 25 ± 1 °C was investigated using potentiody- namic polarization and immersion tests. Results indicated that the microstructure of fusion zones exhib- ited dendritic structure contained lathy and skeletal d-ferrite. The contents of d-ferrite in the weld zone increased by increasing the number of passes. Therefore, as the number of passes increased, the hardness and corrosion resistance increased. Ó 2013 Elsevier Ltd. All rights reserved. 1. Introduction Stainless steels are an important class of engineering materials that have been used widely in a variety of industries and environ- ments due to their high corrosion and oxidation resistance. They typically contain more than 12% chromium which provides the resistance to corrosion by formation of a tenacious, spontaneously regenerating surface layer of chromium oxide. Other alloying ele- ments such as nickel, manganese and molybdenum are often added regarding to different purposes specific to the application. A small amount of carbon is also present, either intentionally added or as an impurity [1]. Austenitic stainless steels are a group of steels that contain nominally 19% chromium and 9% nickel. As implied by the name, austenite is the predominant microstructural phase in these steels at room temperature. The chemical composi- tion can be varied, for example by lowering the carbon content, and by adding titanium, niobium or tantalum to prevent chromium carbide formation, or by adding molybdenum to increase resis- tance to localized corrosion [2]. This group of steels exhibit a highly attractive combination of high strength, good ductility, excellent corrosion resistance and a reasonable weldability. These properties make austenitic stainless steels as an attractive candidate materi- als for use in a wide range of industrial field such as steam-power generation, automotive engineering, biomedical and dairy indus- try, petrochemical and chemical engineering [3]. Among the many 300 series austenitic stainless steel grades, AISI 304L stainless steels are extensively used in industries due to their superior low temperature toughness and high corrosion resistance. It is esti- mated that the 304 and 304L austenitic stainless steels are used to the extent of more than 20,000 tons a year [4]. Welding is one of the most widely used processes to fabricate stainless steel structures. The welding processes have been used in different applications such as pipelines, automotive exhaust gas systems, chemical industrial equipment repairs [5,6]. However, welding often has potent effects on the microstructure and hence is expected to have strong influence on the mechanical and corro- sion properties of the welded samples. Different metallurgical problems such as micro-segregation, precipitation of secondary phases, presence of porosities, cracking and loss of materials by vaporization are the main challenges of fusion welding processes [5,7]. It is well known that the microstructure of austenitic stain- less steel is mainly composed of austenite under the condition of equilibrium solidification. However, during the non-equilibrium rapid solidification conditions, such as in welding, the high cooling rate will result in incomplete c ? d transformation and small amounts of d-ferrite should be remained unavoidably in the weld microstructure at room temperature. The retained d-ferrite is known to prevent solidification and hot cracking and to improve ductility, toughness and corrosion resistance. However, it is also reported that excess d-ferrite (usually more than 10 vol.%) can decrease the hot workability [8]. Therefore, d-ferrite content in the weldments plays an important role in the welding of austenitic stainless steels which should be controlled. 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.10.064 Corresponding author. Tel.: +98 311 3915735; fax: +98 311 3912752. E-mail address: [email protected] (M. Atapour). Materials and Design 55 (2014) 905–911 Contents lists available at ScienceDirect Materials and Design journal homepage: www.elsevier.com/locate/matdes

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Materials and Design 55 (2014) 905–911

Contents lists available at ScienceDirect

Materials and Design

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

Microstructure and corrosion behavior of multipass gas tungsten arcwelded 304L stainless steel

0261-3069/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.matdes.2013.10.064

⇑ Corresponding author. Tel.: +98 311 3915735; fax: +98 311 3912752.E-mail address: [email protected] (M. Atapour).

G.R. Mirshekari, E. Tavakoli, M. Atapour ⇑, B. SadeghianDepartment of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran

a r t i c l e i n f o

Article history:Received 11 June 2013Accepted 24 October 2013Available online 5 November 2013

Keywords:Gas tungsten arc weldingMicrostructureCorrosion resistance

a b s t r a c t

The purpose of this study was to discuss the effect of single pass and multipass (double and triple pass)gas tungsten arc welding (GTAW) on microstructure, hardness and corrosion behavior of 304L stainlesssteel. In this investigation, 308 stainless steel filler metal was used. Microstructures and hardness of theweldments were investigated using optical microscopy, scanning electron microscopy (SEM), energy dis-persive spectroscopy (EDS), X-ray diffraction analysis (XRD) and Vickers microhardness (HV0.5). A ferrite-scope was also used in the non-destructive evaluation to observe the ferrite content on the weldments.Corrosion behavior of weldments in 1 M H2SO4 solution at 25 ± 1 �C was investigated using potentiody-namic polarization and immersion tests. Results indicated that the microstructure of fusion zones exhib-ited dendritic structure contained lathy and skeletal d-ferrite. The contents of d-ferrite in the weld zoneincreased by increasing the number of passes. Therefore, as the number of passes increased, the hardnessand corrosion resistance increased.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Stainless steels are an important class of engineering materialsthat have been used widely in a variety of industries and environ-ments due to their high corrosion and oxidation resistance. Theytypically contain more than 12% chromium which provides theresistance to corrosion by formation of a tenacious, spontaneouslyregenerating surface layer of chromium oxide. Other alloying ele-ments such as nickel, manganese and molybdenum are oftenadded regarding to different purposes specific to the application.A small amount of carbon is also present, either intentionallyadded or as an impurity [1]. Austenitic stainless steels are a groupof steels that contain nominally 19% chromium and 9% nickel. Asimplied by the name, austenite is the predominant microstructuralphase in these steels at room temperature. The chemical composi-tion can be varied, for example by lowering the carbon content,and by adding titanium, niobium or tantalum to prevent chromiumcarbide formation, or by adding molybdenum to increase resis-tance to localized corrosion [2]. This group of steels exhibit a highlyattractive combination of high strength, good ductility, excellentcorrosion resistance and a reasonable weldability. These propertiesmake austenitic stainless steels as an attractive candidate materi-als for use in a wide range of industrial field such as steam-powergeneration, automotive engineering, biomedical and dairy indus-try, petrochemical and chemical engineering [3]. Among the many

300 series austenitic stainless steel grades, AISI 304L stainlesssteels are extensively used in industries due to their superior lowtemperature toughness and high corrosion resistance. It is esti-mated that the 304 and 304L austenitic stainless steels are usedto the extent of more than 20,000 tons a year [4].

Welding is one of the most widely used processes to fabricatestainless steel structures. The welding processes have been usedin different applications such as pipelines, automotive exhaustgas systems, chemical industrial equipment repairs [5,6]. However,welding often has potent effects on the microstructure and henceis expected to have strong influence on the mechanical and corro-sion properties of the welded samples. Different metallurgicalproblems such as micro-segregation, precipitation of secondaryphases, presence of porosities, cracking and loss of materials byvaporization are the main challenges of fusion welding processes[5,7]. It is well known that the microstructure of austenitic stain-less steel is mainly composed of austenite under the condition ofequilibrium solidification. However, during the non-equilibriumrapid solidification conditions, such as in welding, the high coolingrate will result in incomplete c ? d transformation and smallamounts of d-ferrite should be remained unavoidably in the weldmicrostructure at room temperature. The retained d-ferrite isknown to prevent solidification and hot cracking and to improveductility, toughness and corrosion resistance. However, it is alsoreported that excess d-ferrite (usually more than 10 vol.%) candecrease the hot workability [8]. Therefore, d-ferrite content inthe weldments plays an important role in the welding of austeniticstainless steels which should be controlled.

Fig. 1. Schematic illustration of the weldment specimen.

906 G.R. Mirshekari et al. / Materials and Design 55 (2014) 905–911

The most important property of stainless steels, and the reasonfor their existence and widespread use, is their corrosion resis-tance. The corrosion resistance of stainless steels depends on thechemical composition, structure and homogeneity of their passivefilm [9]. It is highly probable that the welding processes can influ-ence the passivity due to thermal cycles and compositionalchanges associated with the welding. Stainless steel weldmentscorrosion has been the subject of considerable research. Intergran-ular corrosion due to sensitization is one of the most commonproblem encountered in stainless steel weldments. This is a wellknown phenomenon that is found to be due to precipitation ofcarbides and other intermetallic phases at grain boundaries duringheating in the sensitization temperature range, typically 600–850 �C. Fusion welding processes can cause local changes in com-position of the welded material which can alter the stability ofthe passive layer and its corrosion behavior [10,11].

Among the various welding techniques, gas tungsten arc weld-ing (GTAW) is the most reliable method for welding stainless steelsdue to its significant advantages like relative cleanliness, ability ofwelding complicated shapes with large and small dimensions andavailability of a wide range of materials. Also, this method is verycost-effective and techno-economic solution for joining of materi-als due to its low price equipments [12]. However, there are scantyinformation regarding to the influence of the number of passes onthe corrosion behavior of stainless steels after GTAW.

The aim of this investigation is to study the corrosion behaviorof the multipass welded 304L stainless steel using 308 filler metals.In addition, the microstructural aspects are also studied.

2. Experimental details

2.1. Materials

The base material used in the present study was in the form ofAISI 304L stainless steel (304L SS) plates with dimensions of200 mm � 100 mm � 6 mm which cut from a rolled sheet. The AISI308 stainless steel (308 SS) solid electrode of 2.4 mm diameter wasselected as the filler metal. The chemical compositions of the basemetal and the filler metal are presented in Table 1.

2.2. Welding procedure

In the present work, plates of AISI 304L SS were butt-jointedusing AISI 308 SS filler in three different conditions (one to threepasses) by GTAW process. Fig. 1 shows the schematic diagram ofthe weldment specimen. Before welding all the sheets werecleaned mechanically and chemically by acetone in order to re-move any source of contaminations like rust, dust, oil, etc. One,two and three passes of welding were performed on the sampleswith the conditions mentioned in Table 2. The interpass tempera-ture of around 150 �C was maintained for second and third passes.No preheat and post heat treatment was carried out on the speci-mens. During and after the welding, joints were visually inspectedfor their quality and it was ensured that all weld beads possessedgood geometrical consistency and were free from visible defectslike surface porosity, blow holes etc. In addition, argon purge at arate of 10 l/min (before welding: 10 s and after welding: 4 s) was

Table 1Chemical compositions of the base metal and the filler metal (in wt%).

Material Fe C Si Mn

Base (304L SS) Balance 0.03 0.35 2.0Filler (ER 308 SS) Balance 0.07 1.0 1.57

used during welding to prevent penetration of N2, O2 and H2 intothe weld zone which could harm properties of the joint.

2.3. Microstructural examination

The cross-section of the test specimens were mounted andmechanically ground to 1200 mesh on SiC papers and finallypolished on the cloth using a suspension of alumina powder. Anaqueous solution HCl:HNO3 (3:1) was used for etching the speci-mens. The microstructure of the weldments was investigated byoptical microscopy and scanning electron microscopy (SEM) mod-el-PHILIPS XL30 equipped with an energy dispersive spectroscopy(EDS) detector. The phase analysis was conducted by X-ray diffrac-tion analysis (XRD) using Philips X Pert-MPD System with Cu Karadiation and 40 kV and 30 mA operating conditions. A microhard-ness tester equipped with Vickers pyramid indenter was used formicrohardness measurements along the longitudinal centerlineof the welds. A 500 g load was applied on the indenter for 10s.For the sake of convenience, three different welds, single to triplepass welds, are denoted by No. 1, No. 2 and No. 3 specimensrespectively.

A ferritescope (MP30E-Fischer) was used in the non-destructiveevaluation to observe the ferrite content on the weldments interms of the ferrite number. According to the various cooling ratesoccurred in different regions of the weld zone and thus affect fer-rite contents, all these tests were done at the center of the welds.

2.4. Electrochemical and immersion testing

For electrochemical studies, samples of the base metal and dif-ferent weld metals were cut with exposed area of 1 cm2. Beforeelectrochemical tests, all samples were mechanically ground usingSiC papers up to 1200 grit and then polished by alumina powder toguarantee consistent surface roughness. The samples were de-greased in acetone, followed by cleaning in distilled water anddried thoroughly in air. To avoid the crevice corrosion during thecorrosion tests, the interface between sample and mounting mate-rial was covered by lacquer.

Corrosion parameters were evaluated by potentiodynamicpolarization tests in 1 M H2SO4 solution at 25 ± 1 �C and aeratedconditions using the advanced electrochemical system, modelPARSTAT 2273 according to ASTM: G5 standard. A standard threeelectrode cell was used with calomel electrode (SCE) and graphiterods as the reference and counter electrodes, respectively. The testswere started after a steady open circuit potential that was attainedat about 30 min. Potentiodynamic polarization measurements

Cr Ni Mo S P

18.02 8.0 0.18 0.004 0.0118.15 10.01 – 0.03 0.04

Table 2Welding parameters used for fabricating butt welded joints.

Filler Filler diameter (mm) Electrode Electrode diameter (mm) Welding speed (mm/min) Current (A) Voltage (V)

308 SS 2.4 W�2%Th 2.4 150 59 12

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were conducted from 250 mV below the open-circuit potential at ascan rate of 1 mV/s. Each test was repeated three times to ensurereproducibility.

Immersion corrosion tests in 1 M H2SO4 solution at 25 ± 1 �Cwere performed on all weldment samples for 48 h to determinethe morphology of the corroded surfaces. Before testing, all sam-ples were mounted and polished, then cleaned in acetone and dis-tilled water. After immersion, the samples were cleaned in distilledwater and ethyl alcohol. Finally, the samples were dried in air andexamined using SEM.

Fig. 3. Typical dendritic structure of the weld zone – the backscatter SEMmicrographs of: (a) the cross-sectional view of the dendrite arms accompanied bythe interdendritic structure, (b) the location of the different phases present in themicrostructure.

Table 3EDS analysis of different phases in the microstructure weld zone (in wt%).

Region Fe Ni Cr

1 61.95 15.78 18.022 59.55 10.46 24.52

3. Results and discussion

3.1. Microstructural and metallographic analysis

The XRD patterns of the 304L SS base metal and three differentwelds are shown in Fig. 2. As can be seen, only austenite peakswere observed in the base metal. But, XRD is not capable of detect-ing phase under 5% wt. Thus, very small amount of d-ferrite may bepresent in the structure that cannot be detected with XRD analysis.In addition, the weld zone contained austenite and d-ferrite phases.During the cooling process, the primary d-ferrite solidifies in thefusion zone and then the d ? c transformation takes place. Sincethe d ? c transformation is a diffusion controlled process, the fastcooling rate in GTAW process does not offer sufficient time to com-plete this phase transformation. As a result, the primary d-ferrite isretained in weld zone [13]. A small amount of d-ferrite is necessaryto avoid the problem of hot cracking during weld solidification.Fully austenitic weld deposits are susceptible to microfissuringduring cooling upon solidification [14].

Fig. 3(a) shows a typical backscatter SEM image of the dendriticstructure present in the weld zone. The backscatter mode in thehigher magnification image shown in Fig. 3(b) reveals two distinctregions in the microstructure denoted by 1 and 2. Region 1 belongsto the dendrite phase while region 2 belongs to the interdendriticphase. EDS analysis data for these two regions are given in Table 3.The EDS data for microstructural regions are being qualitativelydealt with in order to reveal the difference in concentrations ofthe metallic elements between various phases. According to theXRD analysis and EDS data listed in Table 3, it can be concludedthat the dendrite phase (region 1), is austenite (nickel-enrichedphase) and the interdendritic phase (region 2) is d-ferrite (chro-

Fig. 2. X-ray diffraction patterns from the base metal and single to triple passwelds.

mium-enriched phase). Moreover, according to the SEM imagemode, the brighter interdendritic phase has higher chromium con-tent than the dendrite phase. It is resulted since, on the SEMimages taken in the backscatter mode, a chemical compositioncontrast may be observed between different phases based on theatomic number of the major elements in each phase (i.e. the higheratomic number brighter the appearance).

Fig. 4 shows the variation of d-ferrite morphology in the fusionzone. It is clear that two different kinds of lathy d-ferrite and skel-etal d-ferrite were formed in the austenite matrix. Similar d-ferritemorphologies were observed in the fusion zone of all specimens.This phenomenon is attributed to the various thermal cycles andcooling rates in the weld zone. Different cooling rates in multipasswelding have dictated different types of ferrite transformation andtherefore have changed the morphology of ferrite within the weld.

Fig. 4. Typical morphologies evolution of d-ferrite in the fusion zone: (a) lathy d-ferrite and (b) skeletal d-ferrite.

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Where the cooling rates were moderate, a skeletal d-ferrite mor-phology has been formed. This is a consequence of the advanceof the austenite consuming the ferrite until the ferrite is suffi-ciently enriched in ferrite-promoting elements, such as chromium,and depleted in austenite-promoting elements, such as nickel, thatis stable at lower temperatures where diffusion is limited. But,where the cooling rates were high, a lathy d-ferrite morphologyhas been formed. The lathy morphology has been formed due to re-stricted diffusion during the ferrite–austenite transformation.Decreasing the diffusion distances increases the efficient of thetransformation to proceed and therefore decreases spaced laths[14].

Optical micrographs showing the microstructures of heat-af-fected zone (HAZ), fusion boundary and weld zone for No. 1, No.2 and No. 3 specimens are illustrated in Fig. 5. These results showthat the denderitic arm spacing and the HAZ grain size have beenincreased with decreasing number of passes. This behavior can beattributed to the thermal cycling and recrystallization occurreddue to the multipass welding [15]. As illustrated in Fig. 5, an epi-taxial growth occurred in all specimens. This kind of growth is ben-eficial to the welds, since it incorporates the weld zone to thegrains from the base metal. This therefore prevents the weld zone– HAZ interface becoming a region with a concentration of stresses.Similarities in crystal structure and chemical composition betweenthe weld zone and the base metal are the main characteristics ofthis growth [16].

Determination of the ferrite content is an important issue be-cause of its effects on the hot cracking and joint properties. The fer-rite numbers of the welds for No. 1, No. 2 and No. 3 specimenswere 6.3, 7 and 8.2 respectively. This increasing is a result of thevariation of thermal cycling and cooling rates by increasing the

number of passes. Since the d ? c transformation is a diffusioncontrolled process, high cooling rates does not offer sufficient timeto complete the phase transformation. As a result, higher coolingrates causes more retained d-ferrite in the weld zone. It has beenrecommended that 5–10% d-ferrite content is required to preventsolidification and hot cracking and to improve ductility, toughnessand corrosion resistance [1], According to the ferrite number of dif-ferent weld zones, it is concluded that all the specimens are undera secure range against these problems.

3.2. Microhardness

Microhardness measurements were taken in the longitudinaldirection i.e. parallel to the base plate surface. As shown in Fig. 6,despite some instabilities on the measured values, a clear trendof continuously increasing hardness in order of weld zone, HAZand base metal can be observed in all No. 1, No. 2 and No. 3 spec-imens. Maximum hardness values in the weld zone and HAZ wereobtained for No. 3 specimen and then decreased in No. 2 and No. 1respectively. These behaviors are due to the higher d-ferrite in theweld zone and finer grain size in the HAZ by increasing the numberof passes. It is widely known that the presence of d-ferrite can im-prove the mechanical strength [15]. Moreover, minimum hardnessvalues were obtained for weld zone of each specimen. In the weldzone, the heat from the process causes annealing and recovery totake place, leading to a drop in hardness. Although a differencein grain structures in weld zone, no apparent changes in hardnesswere disclosed. Hardness values then gradually increased in theHAZ from near the fusion line to near the base metal. This trendof hardness in the HAZ is due to slow cooling rate of the area nearthe fusion line and subsequent grain growth, In contrast, the areanear the base metal experienced higher cooling rate due to steeperthermal gradients and consequently exhibited a fine grainedmicrostructure [17]. The base metal showed the highest values ofhardness due to more chemical and microstructural homogeneity.

3.3. Corrosion behavior

The potentiodynamic polarization curves of the base metal andweld zones for No. 1, No. 2 and No. 3 specimens in 1 M H2SO4 solu-tion at 25 ± 1 �C are shown in Fig. 7. The corrosion current densities(icorr), passive current densities (ip) and corrosion potentials (Ecorr)of various specimens were determined from the potentiodynamicpolarization curves (Tafel extrapolation method) and summarizedin Table 4. The results show that as the number of passes increased,the corrosion current densities decreased. The corrosion currentdensity decreased from 8.75 lA/cm2 for No. 1 specimen to2.74 lA/cm2 for No. 2 and 1.02 lA/cm2 for No. 3 respectively. Thisbehavior can be explained according to the chromium content.Since the chromium content is higher in the d-ferrite than in theaustenite phase, as the content of d-ferrite phase in the microstruc-ture of welds increased, the corrosion resistance improved. On theother hand, redistribution and/or removal of MnS inclusions whichare the preferred sites for corrosion initiation may be another rea-son for improving the corrosion resistance in No. 3 specimen. Sincesolubility of sulfur in the d-ferrite is higher than that in the austen-ite, it is suggested that the d-ferrite formation reduces the avail-ability for the formation of sulfur conclusions. Therefore, theformation of d-ferrite was considered to be the key factor contrib-uting to the improvement in corrosion resistance [18]. However,the presence of d-ferrite with higher chromium content may havedeleterious effect on the corrosion resistance due to the potentialdifference between the d-ferrite and austenite phases [19], but itseems that this factor had less influence on the corrosion resistancethan the others. Moreover, corrosion properties of the welds alsoinclude grain boundary effects [20–22]. The impurities enrich at

Fig. 5. Optical micrographs showing the microstructure of HAZ and weld zone for: (a, b) No. 1, (c, d) No. 2 and (e, f) No. 3 specimens respectively.

Fig. 6. Microhardness profile showing hardness of different zones of the weldments.

G.R. Mirshekari et al. / Materials and Design 55 (2014) 905–911 909

grain boundaries during solidification of weld zone, making grainboundaries more likely to be active sites for corrosion. The largerarea of grain boundary produced by small grain size can dilutethe impurities and make the grain boundaries less attackable[23], so that the No. 3 specimen becomes the most corrosionresistant. In addition, it should be emphasized that the corrosionbehavior of d-ferrite depends on the corrosive environments. Incorrosive environments which chlorides are present, the Cl� ions

preferentially attack the ferrite phases and lead to corrosion [19].Furthermore, it has been reported that in reducing environments,d-ferrite was also corroded selectively [10]. However in this inves-tigation, the d-ferrite dissolves less than the austenite owing to thehigher chromium content of d-ferrite (Fig. 3 and Table 3). It seemsthat by increasing the number of passes and d-ferrite subsequently,the corrosion resistance of the welds has been improved in thisenvironment. It should be noted that the preheating of the weld

910 G.R. Mirshekari et al. / Materials and Design 55 (2014) 905–911

due to the previous passes can encourage the diffusion and im-prove the migration of elements to their relative phases.

The formation of passive film is one of the main properties ofstainless steels. However, the compositional heterogeneity duringweld metal solidification may significantly alter the stability ofthe passive film and then the corrosion behavior of the weld metal.According to Fig. 7, the passive current density of No. 1 and No. 2specimens was significantly higher as compared to the No. 3specimen. However, No. 3 specimen exhibited similar passive

Fig. 7. Potentiodynamic polarization curves of the base metal and different weldsin 1 M H2SO4 solution at 25 ± 1 �C.

Table 4Electrochemical corrosion parameters obtained from potentiodynamic polarizationcurves for various samples in 1 M H2SO4 solution at 25 ± 1 �C. Average data arepresented. Data are extracted according to Tafel extrapolation method.

Specimen Ecorr (mV) icorr (lA/cm2) ip (lA/cm2)

Base metal �232 0.83 2.85No. 1 �379 8.75 48.5No. 2 �266 2.74 6.54No. 3 �246 1.02 1.94

Fig. 8. SEM images showing the microstructure of HAZ after immersion in 1 M H2SO4 solu

current density to the base metal. It seems the segregation ofalloying elements decreased for No. 3 because of preheating dueto first and second passes before third pass. In other words, somehomogenization takes place in the No. 3 specimen that improvesthe passivity hence its corrosion behavior.

As corrosion standpoint, the weld metal in this investigationcan be compared to the duplex stainless steels. However, commonduplex stainless steels are generally overalloyed in chromium andmolybdenum. Segregation is not a problem for ferrite phase be-cause of rapid high- temperature diffusion in the bcc ferrite lattice.In contrast, segregation becomes much more significant in the aus-tenite phase. Therefore, the welding practice for duplex stainlesssteels must be accurately specified and closely monitored [24–26].

As illustrated in Fig. 7 and Table 4, the base metal exhibitedhigher corrosion resistance with lower corrosion current density(0.83 lA/cm2) in comparison with the welds. Also, the base metalhas higher corrosion potential (�232 mV) than the welds. Thus,when the base metal and weld are adjacently placed together inthe electrolyte, the weld is the anodic part of the galvanic couplebetween base and weld metals. Since the welds are more sensitiveto corrosion, it is preferable to choose a situation where the weldhas the corrosion potential near the base metal. Accordingly, No.3 specimen is desirable because of its very similar corrosion poten-tial to the base metal.

The HAZ of austenitic stainless steels can also be susceptible tothe intergranular corrosion due to precipitation of chromium car-bides, especially Cr23C6, at grain boundaries, which is called sensi-tization. More chromium at grain boundaries has been consumedto form larger chromium carbides, making grain boundaries lesscorrosion resistant. The precipitation temperature range of the car-bides is about 600–850 �C. Chromium carbides dissolve at abovethis temperature range [1]. Fig. 8 shows the microstructure ofHAZ in all specimens after immersion test in 1 M H2SO4 solutionat 25 ± 1 �C for 48 h. As can be seen, intergranular corrosion inHAZ decreased with increasing the number of weld passes. The re-sults indicate that as the number of passes increased, the corrosionattack of grain boundaries decreased. This behavior is due toincreasing the retention time above sensitization temperature

tion at 25 ± 1 �C for 48 h for: (a) No. 1, (b) No. 2 and (c) No. 3 specimens respectively.

G.R. Mirshekari et al. / Materials and Design 55 (2014) 905–911 911

range which promotes the solubilization of the chromium carbidesin the grain matrix and therefore avoids the sensitization.

4. Conclusions

The microstructure, hardness and corrosion resistance evolu-tions of the weldments were investigated during single pass andmultipass (double and triple pass) GTAW process of 304L SS. Theimportant results can be summarized as follows:

(1) All welds were essentially austenitic with the presence of asmall amount of d-ferrite. The microstructure of the weldswas dendritic which d-ferrite phase placed in interdendriticregions. The morphology of d-ferrite showed the lathy andskeletal d-ferrite formed in the austenite matrix. In addition,the content of d-ferrite in the weld zone was maximum forthe triple pass welded specimen and decreased in doublepass and single pass welded specimens respectively.

(2) The hardness values from the weld zone toward the basemetal increased in all weldments while the maximum val-ues obtained for three passes welded specimen due to theincrease in the d-ferrite content and grain refinement.

(3) Potentiodynamic polarization curves and immersion testsrevealed that as the number of passes increased, the corro-sion resistance also increased. Additionally, three passeswelded specimen possessed the highest corrosion resistancecompared with the other welded specimens.

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