Corrosion Science 51 (2009) 2316–2323

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Corrosion Science

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Intergranular corrosion susceptibility in supermartensitic stainless steel weldments

J.M. Aquino *, C.A. Della Rovere, S.E. KuriSão Carlos Federal University (UFSCar), Materials Engineering Department, Rodovia Washington Luís, km 235, CEP 13565-905, São Carlos, SP, Brazil

a r t i c l e i n f o

Article history:Received 31 March 2009Accepted 8 June 2009Available online 14 June 2009

Keywords:C. Intergranular corrosionC. Pitting corrosionStainless steelWelding

0010-938X/$ - see front matter � 2009 Elsevier Ltd. Adoi:10.1016/j.corsci.2009.06.009

* Corresponding author. Tel.: +55 16 33518506; faxE-mail address: dsek@power.ufscar.br (J.M. Aquino

a b s t r a c t

The intergranular corrosion susceptibility in supermartensitic stainless steel (SMSS) weldments wasinvestigated by the double loop – electrochemical potentiokinetic reactivation (DL-EPR) techniquethrough the degree of sensitization (DOS). The results showed that the DOS decreased from the basemetal (BM) to the weld metal (WM). The heat affected zone (HAZ) presented lower levels of DOS, despiteof its complex precipitation mechanism along the HAZ length. Chromium carbide precipitate redissolu-tion is likely to occur due to the attained temperature at certain regions of the HAZ during the electronbeam welding (EBW). Scanning electron microscopy (SEM) images showed preferential oxidation sites inthe BM microstructure.

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1. Introduction

Intergranular corrosion is a selective process that occurs in sen-sitized regions of stainless steels as a result of inadequate heattreatments, welding, or high-temperature service [1]. In the classi-cal sensitization process there is the formation of chromium richprecipitates at grain boundaries, with the subsequent impoverish-ment of that element in the adjacent matrix. Consequently, passivefilms formed over those depleted regions are not stable, thus, itleads to a more susceptible region to a corrosion attack. In that as-pect, intergranular corrosion could resemble a galvanic cell inwhich the grains are the cathodic area, and the corresponding grainboundaries are the anodic one; which results in a high cathodicarea in relation to an anodic area [2]. Martensite induced sensitiza-tion is another type of sensitized region susceptible to an intra-granular corrosion attack [1,3]. This process occurs due to thechromium carbide precipitation in the martensitic lath boundaries[3] which is more intensive in tempered than in quenched mar-tensite [4]. Other metallurgical phases may contribute to influencethe corrosion resistance. According to the literature [4] austenitepromotes carbon and nitrogen dissolution, having a consequentreduction in the chromium and molybdenum precipitates. Thus,pit potential demonstrated to be dependent on the austenite con-tent; which gave noble potentials [5]. Otherwise, the d-ferritephase presented in low carbon 13%Cr steels deteriorates the corro-sion resistance [6] due to the higher carbide precipitation aroundthat phase, that is occasioned by its low carbon solubility [7]. Othercorrosion and mechanical problems are associated with that phase[8,9].

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: +55 16 33615404.).

In welded joints, chromium and molybdenum carbonitride pre-cipitations in the heat affected zone (HAZ), particularly in thesupermartensitic stainless steel (SMSS), are responsible for its sus-ceptibility to a corrosion attack. Those precipitations mainly occurat the prior austenite grain boundaries and also at the martensite/d-ferrite phase boundaries during a multipass welding, with anadditional intensification promoted by the post weld heat treat-ments (PWHT) [10]. Carbide precipitates occur due to the combi-nation of carbon matrix saturation and tempering effect of thesubsequent welding passes. That process is not expected to occurin the cap layer or in the single pass welded joints [10]. Intergran-ular stress corrosion cracking (IGSCC), which was not expected inthe HAZ of the SMSS, can occur due to small amounts of Cr-carbideprecipitates at the prior austenite grain boundaries [11]. Accordingto the literature [12,13], the HAZ is the most susceptible region to acrack initiation, due to its sensitization in weldments made by con-ventional welding processes.

Additionally, the welding technique exerts influence on the cor-rosion behavior of the weldment. The high power density pro-cesses are interesting due to their high heat input; which isconfined in a small region of the workpiece, and also to their highcooling rates. These characteristics promote the dissolution of Cr-carbide precipitates as well as their suppression. That precipitatesuppression is caused by very short time periods in the tempera-ture precipitation zone [14]. Corrosion resistance and mechanicalproperties, when comparing the high power density to the conven-tional processes [14,15], can lead to better welded joints.

Pitting corrosion is also likely to occur at chromium depletedregions [5] as well as near the MnS inclusions [16–21]. However,a direct relationship was not established. According to the litera-ture [22], the HAZ of a high power density process weldment isnot the most susceptible region to the pitting corrosion due to

Table 2Electron beam welding parameters for the top pass.

Parameters Conditions

Vacuum 1 mbarWelding voltage 60 kVWorking distance 50 mmWelding speed 7 mm s�1

Root opening 0.5 mmWire feed rate 3.78 m min�1

Welding current 130 mAHeat input 1.1 kJ mm�1

J.M. Aquino et al. / Corrosion Science 51 (2009) 2316–2323 2317

the redissolution process of the Cr-carbide precipitates. Further-more, the HAZ microstructure does not seem to influence the pit-ting corrosion [23].

The double loop electrochemical potentiokinetic reactivation(DL-EPR) is a corrosion testing method characterized by a rapid,quantitative, and non-destructive test to measure the degree ofsensitization [24]. This technique was primarily developed foraustenitic stainless steel using standard practices. However, inter-granular corrosion concerning the martensitic stainless steel (MSS)is a potential problem due to its use in the tempered condition forhardness adjustment. As the conventional MSS contains high levelsof carbon, it is not regularly welded. On the other hand, with thedevelopment of new weldable supermartensitic stainless steel(SMSS), which has lower carbon content, there is the need to inves-tigate the behavior of those welded joints regarding the sensitiza-tion. In order to combine less susceptible welded joints, electronbeam welding is inserted, considering intergranular and pittingcorrosion. That process links high welding velocities to narrowheat affected zones of a few millimeters, with single passes, andno PWHT. These characteristics enable a more resistant weldment,considering precipitation as well as corrosion.

The aim of this work was to measure, comparatively, the degreeof sensitization (DOS) of the SMSS weldments, through the DL-EPRtechnique.

2. Experimental

Electrochemical measurements were done in samples of thebase metal (BM), the heat affected zone (HAZ), and the weld metal(WM); in two classes of supermartensitic stainless steel that werewelded by an electron beam in a chamber of low vacuum.

The full penetration butt weld was done between two 20 mmthick plates, which were previously hot-rolled and tempered at600 �C for 10 min. The welding was done by using a matching fillerwire metal, in two passes, without PWHT. The chemical composi-tion of the plates, the filler wire, and the welding operation condi-tions are showed in Tables 1 and 2, respectively.

The 1 mm thick samples were extracted by electroerosion, fromthe inner part of the weldment top. Fig. 1(a–c) shows a schematicrepresentation of the samples’ extraction. Each extraction step wascarried after etching the weldment sample with a villela’s reagent,through the identification of the characteristic microstructures(BM, HAZ, and WM), as well as through an optical microscopicanalysis. In order to attempt disregarding any galvanic influence,the samples’ electrochemical investigation was carried separatelyin the BM and in the WM regions. A clear HAZ sample cutoutwas not possible to be obtained due to its small size. Fig. 1(d)shows that the sample also presented BM and WM region residues.

X-ray diffraction measurements were done separately in theBM, HAZ, and WM samples for identifying and quantifying themetallurgical phases. The austenite phase quantification (volumet-ric fraction, %) was carried according to the standard practiceASTM-E975 [25]. The X-ray beam was focused in the middle ofall samples. That was particularly important when consideringthe HAZ samples. The analysis was a 2h–h with a continuoussweeping rate equal to 2 degrees a minute, ranging from 5 degreesto 120 degrees. During the X-ray measurements the samples wererotationed.

Table 1Chemical composition (mass%) of steels and filler metal.

Material C Si Mn P S Cr

A 0.02 0.3 0.9 0.03 0.004 12.6B 0.007 0.07 1.6 0.007 0.003 11.4Filler 0.012 0.45 0.65 0.005 0.005 12.3

The electrochemical testing samples were mounted in polyesterresin to avoid the presence of crevices after the electric contact.The exposed area was 0.5 cm2. Before the polarization measure-ments, the samples were wet-grinded on 600 silicon carbide(SiC) paper, washed in distilled water, and immersed in the elec-trochemical cell.

The test solution was 0.5 mol L�1 H2SO4 with 0.01 mol L�1

KSCN, reagent-grade, in distilled water. The electrolyte was natu-rally aerated and the temperature was held at 25 �C. A conven-tional electrochemical cell composed of a platinum counterelectrode and a satured calomel reference electrode (SCE), con-nected to a Solartron SI 1287A potentiostat, was used.

The DL-EPR polarization curves were obtained after two steps:First, the working electrode was subjected to open circuit condi-tions, until a steady state potential (Ecorr) was reached. This wasaccomplished in 25 min. Then, an anodic potentiodynamic sweep-ing rate of 1.67 mV s�1, from �100 mV/Ecorr to +600 mVSCE, wasimposed. At +600 mV, the potential scanning was reversed backto �100 mV/Ecorr. At least four curves for each sample region weredone to obtain good reproducibility. The test results were ex-pressed in the current densities’ ratio, iR/ia, which was used to eval-uate the DOS. The iR term is the reactivation current density(maximum current density in the reverse scan), and the ia termis the activation current density (maximum current density inthe anodic scan).

The DOS was compared to the pit potential results, which weredone in a separate work [22]. The used weldment samples werethe same in both works. The pit potential determination was donein a 3.56 mass% (reagent-grade) sodium chloride solution.

3. Results and discussion

Fig. 2(a–c) shows some of the initial microstructures of the BM,the HAZ, and the WM. Due to their similar microstructural appear-ance in the weldment regions of both steels, there were chosenonly three micrographs. Tempered martensite is the main metal-lurgical phase presented in the BM as a result of the temperingheat treatment. The HAZ microstructure also exhibited temperedmartensite as a consequence of the thermal heat gradient. Thatprocess enabled distinct levels of tempering in the HAZ, as a conse-quence of the different temperatures reached. Additionally, theHAZ was characterized by a heterogeneous microstructure (rela-tive to grain size). Quenched martensite is the WM characteristicphase that is formed due to the high cooling rates attained. Re-tained austenite was presented in the BM (A steel = 11.1% and Bsteel = 39.0%), in the HAZ (A steel = 5.5% and B steel = 4.3%), and

Ni Mo Ti V Cu O N

5.1 1.8 0.01 0.05 0.3 0.01 0.016.1 2.6 0.02 0.05 0.5 0.01 0.016.4 2.6 – – – 0.008 0.01

Fig. 2. Optical micrograph of the weldment: (a) B steel BM, (b) A steel HAZ (next tothe fusion line), and (c) A steel WM. The arrows indicate the d-ferrite phase.

Fig. 1. Sample extraction schematic representation of an electron beam weldment top pass: (a) welded joint, (b) weldment slice extracted by electroerosion, (c) testingsamples having dimensions of 10 mm � 5 mm � 1 mm, and (d) the HAZ sample was composed of the BM and the WM region. Weldment regions 1, 2, and 3 refer to the basemetal (BM), heat affected zone (HAZ), and weld metal (WM). The asterisk indicates where the electrochemical testings were done.

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in the WM (A steel = 26.4% and B steel = 18.5%). According to theliterature [7], the austenitic phase in low carbon MSS can only bevisualized through transmission electron microscopy (TEM). d-Fer-rite was detected in the BM, and also next to the fusion line in theHAZ. Its morphological appearance resembles dark stringers asindicated by the arrows in Fig. 2(a and b).

Fig. 3(a and b) shows representative curves of the DL-EPR testfor the BM region of the A and B steels. The activation (anodicdirection) and reactivation (cathodic direction) current densitiesare similar, and very high. That indicates a high sensitized material(iR/ia = 0.589 and 0.643 for the A and B steel, respectively). Scan-ning electron microscopy (SEM) images of Fig. 4(a and b), revealintergranular and intragranular corrosion attacks, occasioned bythe chromium carbide (Cr-carbide) precipitates, and possibly otherprecipitates like Fe2Mo [26]. That process occurred during the BMtempering heat treatment.

Fig. 3 also shows the presence of two anodic peaks in the acti-vation process of both steels. In order to separate and investigatethe origin of these peaks, a potentiokinetic sweeping rate at0.67 mV s�1 was only conducted in the anodic direction. Fig. 5shows that it was not possible to separate, completely, the twoanodic peaks of the B steel. According to the literature [3,27–30],there are some interpretations for the so-called second anodic cur-rent maximum (SACM), which are: (1) nickel enrichment on thesurface; (2) adsorbed hydrogen oxidation; (3) Fe2+ ions effect; (4)chromium depleted zone effect, and (5) microstructural and com-positional effects. Fig. 6(a and b) shows SEM images obtained afterthe potential sweeping rate interruption at E1 = �200 mVSCE, and atE2 = 50 mVSCE, respectively, which are indicated by the arrows in

Fig. 3. DL-EPR curves for the A and B steel BMs. The potential scanning direction isindicated herein (passive vertex potential at 600 mV). The ia and iR current densitiesare also indicated.

Fig. 4. SEM micrographs of the BM samples after a DL-EPR testing: (a) A steel, and(b) B steel. The samples were previously polished in alumina 1 lm.

Fig. 5. B steel BM anodic polarization at 0.67 mV/s from �100 mVOCP to+600 mVSCE.

Fig. 6. SEM micrographs of the B steel BM after a DL-EPR testing up to: (a)E1 = �200 mVSCE, and (b) E2 = 50 mVSCE. Microstructure revealed after a singlepolarization on the anodic direction, for the same testing sample. The sample waspreviously polished in alumina 1 lm.

Fig. 7. B steel BM cathodic polarization at 0.67 mV/s from +600 mVSCE to�100 mVOCP.

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Fig. 5. The micrographs revealed a generalized corrosion attackover the microstructures, without a significant difference betweenthem. The expected chromium depleted zone effect for the SACMwas not confirmed in the SMSS, according to the used experimentaltechniques. The same occurrence was observed on the A steel, butwithout any conclusion.

During the reactivation process on Fig. 3, the B steel BM alsopresented two reactivation maximums. The separation and the

investigation of these processes were conducted through a voltagescan only in the cathodic direction, from +600 mVSCE to �250 mV/Ecorr, after the open circuit conditions were reached. Fig. 7 showsthe resulting polarization curve. The arrows at E3 = �100 mVSCE,and at E4 = �190 mVSCE, indicate the potential sweeping rate inter-ruption in order to obtain SEM images. Fig. 8(a) shows that the firstreactivation peak (E3 = �100 mVSCE) corresponds to the martensite

Fig. 8. SEM micrographs of the B steel BM after a DL-EPR testing up to: (a)E3 = �100 mVSCE, and (b) E4 = �190 mVSCE. Microstructure revealed after a singlepolarization on the cathodic direction, for the same testing sample. The sample waspreviously polished in alumina 1 lm.

Fig. 9. DL-EPR curves for the A and B steel HAZs. The potential scanning direction isindicated herein (passive vertex potential at 600 mV).

Fig. 10. SEM micrographs of the HAZ samples next to the BM (2 mm from thefusion line) after a DL-EPR testing: (a) A steel, and (b) B steel. The samples werepreviously polished in alumina 1 lm.

Fig. 11. SEM micrographs in the middle of the HAZ samples (1 mm from the fusionline) after a DL-EPR testing: (a) A steel, and (b) B steel. The samples were previouslypolished in alumina 1 lm.

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phase contour laths corrosion attack. The second peak(E4 = �190 mVSCE) is due to the entire martensitic matrix corrosionattack (intragranular), as showed in Fig. 8(b). According to the lit-erature [24], the second reactivation peak could correspond to anickel rich phase, which is probably related to the retained austen-ite phase in the SMSS. In addition, the ascribed nickel phase corro-

Fig. 13. DL-EPR curves for the A and B steel WMs. The potential scanning directionis indicated herein (passive vertex potential at 600 mV).

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sion attack at E3 is associated with a chromium content decreasethat caused the potential shift toward negative values. Such reacti-vation occurrence was not observed in the A steel BM sample dueto its lower austenite content. The reactivation splitting occurrenceof the current density peak occurs on the martensite–austenitesteels containing the following [31]: (i) at least 13% of chromium;(ii) alloyed with molybdenum; (iii) and an austenite content higherthan 25% resulting from a suitable tempering heat treatment. Thatoccurrence, on the DL-EPR reactivation curve, was detected instainless steel type 304 [1]; however, having interpretations differ-ent from the ones herein.

The DL-EPR curves of the HAZ samples are representativelyshown in Fig. 9. The resulted curve is a measurement of the entireHAZ from its fusion line to its 2 mm-away point. The B steel degreeof sensitization (iR/ia = 0.005) is much smaller than the A steel one(iR/ia = 0.08). That is because A steel has a higher carbon contentthan B steel, which results in a higher Cr-carbide precipitation, asit will be showed in the micrographs. However, when comparingthe HAZ samples to the BM samples, a much lower DOS is obtainedbecause of the Cr-carbide precipitates possible redissolution at theHAZ specific zones. That process occurred due to the gradient’shigh-temperature in the HAZ, which is proper to the precipitatesredissolution [14,32]. Furthermore, a reduced time for the Cr-car-bide precipitation could also contribute to the DOS lower levelsin that region due to the EBW high cooling rate [14]. Studiesregarding the Cr-carbide precipitate redissolution are in progress.

The second maximum current density in the activation/reacti-vation processes were not observed in the HAZ samples, evenwhen 0.67 mV s�1 sweeping rates were used.

Fig. 12. SEM micrographs of the HAZ samples next to the fusion line after a DL-EPRtesting: (a) A steel, and (b) B steel. The samples were previously polished in alumina1 lm.

SEM images were obtained in three distinct regions within theHAZ samples of both steels, which were: (1) next to the BM(2 mm from the fusion line), (2) in the middle of it (1 mm fromthe fusion line), and (3) next to the fusion line.

Fig. 10(a and b) shows representative images of the HAZ next tothe BM. An outstanding corrosion attack, with a high number ofholes, can be observed on the grain boundaries, and also insidethe A steel grains. The B steel samples are mainly characterized

Fig. 14. SEM micrographs of the WM samples after a DL-EPR testing: (a) A steel, and(b) B steel. The samples were previously polished in alumina 1 lm.

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by an intragranular corrosion attack over the martensitic matrix.This difference is a consequence of the alloying element distinctlevels, mainly carbon [33].

Fig. 11(a and b) shows representative images in the middle ofthe HAZ. Clearly, the intergranular and the intragranular corrosionattacks are decreased in both steels. This is a consequence of thegenerated thermal heat gradient, which enabled the Cr-carbideprecipitate dissolution [14,33]. This process is more prominent inthe martensitic matrix because of the higher carbon diffusion coef-ficient in that phase [7].

Fig. 12(a and b) shows HAZ images next to the fusion line. With-in that region, there is a formation of the d-ferrite phase, which isthe lighter color phase in the micrographs. An intensive corrosionattack can be observed adjacent to that phase, which was revealedby the DL-EPR tests. This process occurred due to the lower carbonsolubility in the d-ferrite phase [7], which resulted in the matrixsupersaturation and the consequent precipitation of the Cr-car-bides that were adjacent to the d-ferrite phase. As the A steel hasa higher carbon content than the B steel, it exhibited a more inten-sive precipitation, which was confirmed by the numerous holes ofthe corrosion attack parallel to the d-ferrite phase. The precipita-tion also occurred in the B steel samples, which were observedthrough the small cavities in Fig. 11(b), in spite of their very smallcarbon content. The HAZ micrographs elucidated the difference be-tween the reactivation current densities of the A and B steels(Fig. 9).

It was not possible to do the HAZ samples’ TEM due to their re-duced number, and to the difficulty of obtaining these samples.

Fig. 15. DOS and pitting potential values within the supermarte

However, new methodologies regarding sample extraction andanalysis [34,35] are being considered.

The WM region DL-EPR curves are shown in Fig. 13. Both steelspresented very similar activation current densities. However, the Asteel samples were more sensitized than the B steel ones due totheir higher reactivation current density (iR/ia = 0.013 and 0.004for the A and B steel, respectively). Nevertheless, the WM regionwas the least sensitized region within the weldment. No evidenceof the second maximum current density was observed in the ano-dic nor in the cathodic scan directions, even after 0.67 mV s�1

sweeping rates were applied. Fig. 14(a and b) shows scanning elec-tron WM images after the DL-EPR test. The A steel presented verysmall corrosion attack cavities, which were almost absent in the Bsteel samples. There is no evidence of preferential precipitationsites, intragranular or intergranular corrosion attacks in thequenched martensite. The carbide precipitates lower levels in theWM region were observed due to the high cooling rates that oc-curred during the welding procedure [36]. This process enabled avery short remaining time in the Cr-carbide precipitation region,and a lower alloying element microsegregation which led to a re-duced intergranular corrosion susceptibility. The high cooling rateeffect also supported the sensitization lower levels in the HAZ dueto the reduced time period within which the material stayed in theCr-carbide precipitation region [14].

The degree of sensitization was compared within different re-gions of the weldment, and is shown in Fig. 15(a and b). The HAZsamples presented levels of sensitization extremely lower whichwere comparable to the WM ones. The HAZ performance concern-

nsitic stainless steel weldments: (a) A steel, and (b) B steel.

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ing the intergranular corrosion was completely different fromwhat is reported in the literature [12,13,35]. However, few re-searches on SMSS, as well as on its weldments that are donethrough high power density processes, are available. The mostimportant point to consider is the welding process, which has dif-ferent characteristics when comparing it to the conventional pro-cess [14,15]. Among those characteristics, the sharp temperaturevariation of the HAZ is the most important because it mainlycauses the Cr-carbide precipitation/dissolution in it.

Fig. 15 also includes the weldment pitting potential values.These results are found in paper [22]. According to the literature[16–20], the presence of chromium depleted regions are relatedto a pit nucleation and growth. As the sensitization increases, thepitting potential decreases, which results in a lower corrosionresistant material.

4. Conclusion

The BM was the most sensitized region within the SMSS weld-ment due to its tempering process, which resulted in its Cr-carbideprecipitation. The WM and the HAZ were the lower sensitized re-gions due to the WM high cooling rates, as well as to the Cr-carbideprecipitate redissolution at certain regions of the HAZ. That wasconfirmed by the SEM images within the HAZ samples.

The HAZ of an electron beam top pass weld is susceptible tosensitization. Moreover, the precipitation within that regionshowed to vary according to the attained temperature duringwelding. The d-ferrite phase (microstructural effect) also influ-enced the HAZ sensitization, because of its low carbon solubility.The austenite content did not seem to influence the sensitizationor the pit potential.

The splitted reactivation peak in the A steel BM samples corre-sponded to two oxidation processes, which were possibly relatedto: (1) the martensitic phase contour laths, and (2) the martensiticmatrix, which may be associated to the austenitic phase content inthe SMSS. The second maximum current density in the activationprocess was not splitted into two peaks. Furthermore, that occur-rence did not seem to be related to the chromium depleted effect.

A correlation between sensitization and pitting potential wasestablished in the SMSS weldments, indicating that the probabilityof pitting corrosion enhanced as the sensitization increased. How-ever, it was not confirmed that chromium depleted regions werethe preferential sites for pit nucleation and growth.

Acknowledgements

The authors thank the financial support provided by CAPES andDr. Celso Roberto Ribeiro for supplying the steel samples.

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