Corrosion degradation of AISI type 304L stainless steel ... · Corrosion degradation of AISI type 304L stainless steel for application in nuclear reprocessing plant ... problems with

ORIGINAL PAPER

Received: 31 December 2014 /Revised: 4 May 2015 /Accepted: 13 May 2015 /Published online: 26 May 2015# Springer-Verlag Berlin Heidelberg 2015

Abstract The degradation of materials involving corrosion inhandling nitric acid in the spent fuel nuclear reprocessing plantis a serious issue. In the present work, the corrosion resistanceof American Iron and Steel Institute (AISI) type 304L stain-less steel (SS) and nitric acid grade (NAG) type 310L SS in 1to 11.5 M HNO3 and boiling 15.65 M HNO3 was evaluated.In both the alloy steels, the open circuit potential and corrosionpotential are shifted to more noble potential with increasingconcentrations. However, the passive current density was notaffected, and the transpassive potential was shifted to higherpotential with increasing concentrations. The corrosion ratemeasured in boiling 15.65 MHNO3 after 240 h shows a muchlower corrosion rate in type 310L SS (∼0.06 ± 0.012 mm/y)then type 304L SS (∼0.18 ± 0.02–0.2 ± 0.001 mm/y). Theseobservations are corroborated with the scanning electron mi-croscope (SEM) morphologies that show severe intergranularcorrosion (IGC) attack in type 304L SS then in type 310L SS.The X-ray photoelectron spectroscopy (XPS) study of thepassive oxide films of both alloy steels shows the presenceof Cr2O3 and SiO2, and the depth profile indicated predomi-nant Si enrichment.

Keywords Corrosion resistance . Nitric acid corrosion .

Stainless steel . Boiling test . XPS . SEM

Introduction

Nitric acid is the primary medium used universally for thewell-proven plutonium–uranium extraction (PUREX) processfor the majority of nuclear fuel reprocessing world over andfor waste storage vessels involving hot concentrated nitricacid [1–3]. Furthermore, nitric acid is the second most indus-trial acid after sulfuric acid [4]. However, nitric acid is strong-ly oxidizing and very corrosive [5–7]. Structurally; austeniticSS are used in many components of reprocessing plant, i.e.,pipework, vessel, tanks, and equipment for handling in therange 60–70 % HNO3 and temperature of operation below80 °C [1, 3, 5]. Austenitic SS readily passivating in nitric acidis attributed due to their ability to form thin, iron/chromiumprotective passive oxide films in highly oxidizing conditionsand, consequently, high corrosion resistance [1, 4, 5].Additionally, the materials for use in the nitric acid applicationhave to be resistant to transpassive corrosion, IGC, and lowgeneral corrosion rate [1, 3, 6]. Therefore, the constructionmaterials used for reprocessing plant equipment have to bechosen carefully considering the nature of medium and con-centrations encounter (i.e., diluted (1–4 M) to concentrated(10–14 M)), room temperature (liquid/solvent extraction),intermediate (warm, waste storage tanks, etc.), boilingtemperature (dissolver, evaporator, etc.), radioactivity,oxidized/redox ions, etc. [1, 3, 5]. The most widely usedaustenitic grades of SS are the low-carbon type 304L SS,stabilized type 310Nb, 321Ti, and type 347Nb SS [1–3, 6,7]. Most austenitic steel alloys resist corrosion over a widerange of temperatures and concentrations up to 60 % HNO3

at its atmospheric boiling points [1, 3, 4]. However, in highlyoxidizing conditions, i.e., higher concentrations (≤8 MHNO3), higher temperatures (≤80 °C and boiling condition),or the presence of oxidizing species, Mn(VII), Ru(IV), Fe(III),Cr(VI), Ce(IV), etc. eluted from spent fuel posed a risk of

* S. [email protected]; [email protected]

1 Corrosion Science and Technology Group, Indira Gandhi Centre forAtomic Research, Kalpakkam 603 102, India

2 Faculty of Engineering, Hokkaido University, Kita-13, Nishi-8,Kita-ku, Sapporo 060 8628, Japan

J Solid State Electrochem (2015) 19:3533–3542DOI 10.1007/s10008-015-2891-y

Corrosion degradation of AISI type 304L stainless steelfor application in nuclear reprocessing plant

S. Ningshen1& M. Sakairi2

material property degradation associated with transpassivecorrosion and IGC attack or even catastrophic plant failure[2, 3, 6]. Furthermore, pitting corrosion, also called tunnelingcorrosion or end-grain attack, occurs depending on the grainorientation (i.e., the corrosion rate increases in the orderplate < side << end grain) of the exposed surface and forgedsteels in highly oxidizing HNO3 conditions [6, 8]. Otherproblems with the corrosion of SS by nitric acid also includevapor-phase corrosion and its variants over very strong acid,and the effects of contaminants and admixtures [2, 3, 9].Therefore, materials for application in corrosive nitric acidenvironment are made of many types of corrosion-resistantmaterials such as austenitic SS; valve/refractory metals, i.e.,Ti, Zr, Hf, Nb and Ta, etc; and its alloys as alternate toaustenitic SS [1, 2, 6, 7, 10]. At lower concentration andtemperature, to avoid transpassive corrosion: addition ofsilicon in austenitic SS is known to increase their corrosionresistance that inhibits the transpassive corrosion and IGCresistance [2, 3, 7, 11, 12]. However, these effects aredependent on the corrosive environment and silicon content,i.e., 2–4 % Si is highly resistant to IGC; however, alloys thatcontain about 0.5–1 % Si corrode faster [12]. Furthermore,most SS, except for certain high-chromium steels, show anaverage corrosion rate of 0.13 mm/y in boiling 65 % HNO3

[1, 6, 10]. These efforts have identified alloys with extra lowlevels of C, Si, P, S, and Mo called the nitric acid-gradestainless steel (NAG SS). The NAG SS [1–3, 6, 7, 10] arecharacterized by (i) control chemical composition of alloyingelements, (ii) microstructures that eliminate IGC sites, and (iii)alloy steels with Si addition and higher Cr that enhancespassive film stability against transpassive dissolution [8, 11].The development of NAG SS similar to AISI types 304LSSand several new proprietary NAG alloy steels used in severalreprocessing plants are described elsewhere [1–3, 7, 10, 11].Thereby, an understanding of how SS interacts withHNO3 in spent fuel plant environments at different con-centrations is essential to enable effective application ofsuch steels in required nuclear environments. In thisstudy, the corrosion behavior of Si-containing type 304LSS and a modified grade of type 310L SS were comparedand investigated in consideration to characteristics ofreprocessing solution (i.e., different concentrations) forspent nuclear fuel. Additionally, immersion test in boilingnitric acid to understand the corrosion acceleration behav-iors was also studied.

Material and experimental methods

Materials

The chemical compositions of type 304L SS (Fe-19.1Cr-9.40–0.54Si-1.60Mn-0.018C-0.04P-0.02S) and type 310L

SS (Fe-25Cr-20Ni-0.012C-0.2Si-0.023P-0.001S) all inweight percent were used in the present work. The spec-imen was wet ground up to 1000 grit using a SiC paperthen polished to mirror finish with diamond paste. Afterthis, the specimens were ultrasonically cleaned in acetoneto remove any surface contaminants, air dried, and thenused immediately for the electrochemical experiment.Similarly, to classify the microstructures, the SS sampleswere electrolytically etched at 1 A cm−2 for 1.5 min using10 % oxalic acid. The etched microstructure of both thealloy steels (Fig. 1) shows the homogeneous austeniticmicrostructure; annealing twins were observed in somegrains and precipitates, or intermetallic phases were notobserved at the grain boundaries and inside grains in boththe alloys.

Corrosion rate measurements

The specimens for boiling tests were wet ground with SiCpaper up to 1000 grit SiC finish. The weight loss measure-ments were conducted by immersion of the specimen into15.65 M (69 %) HNO3 at boiling condition for five 48-hcycles for a total period of 240 h (American Society forTesting and Materials; ASTM A262 practice C Huey test).The solutions were replaced after every 48 h, samples airdried, and the weight loss was measured for each sample be-fore and after such treatment. The corrosion rate was

Fig. 1 Typical optical micrographs observed after etching at 10 % oxalicacid: a type 304L stainless steel and b type 310 stainless steel

3534 J Solid State Electrochem (2015) 19:3533–3542

determined by mass loss method (ASTM G1) using the equa-tion as follows:

Corrosion rate mm=yð Þ ¼ 8:76� 104 �W� �

= A� T � Dð Þð1Þ

where T is time of exposure (h), A is total surface area (cm2),W is weight loss (mg), and D is density (g/cm3).

Electrochemical corrosion measurements

The open circuit potential (OCP) was measured at differ-ent concentrations of 1, 3, 6, 9, and 11.5 M HNO3 in non-stirred condition at room temperature (25 ± 1 °C). Beforethe OCP measurement, the specimens were stabilized for60 min and the OCP was measured as a function of timeup to 60 min.

The potentiodynamic anodic polarization measurementswere carried out at different concentrations of 1, 3, 6, 9, and11.5 M HNO3 at room temperature. The electrochemical testconsisted of three electrodes: reference electrode of Ag/AgClin saturated KCl (0.197 V vs. standard hydrogen electrode(SHE)), the counter electrode platinum (Pt), and the workingelectrode. A solution Luggin bridge was used to provide aconductive path between the working electrode and the refer-ence electrode. The electrode potential was anodicallyscanned at a scan rate of 0.167 mV s−1, and all the electrodepotentials were measured against the Ag/AgCl (in saturatedKCl) reference electrode. A sample surface area of 1 cm2 wasused for all the electrochemical corrosion experiment. Theelectrolyte of all the electrochemical measurements waspurged with purified argon to deaerate the solution, and thepurging was continued till the end of the experiment. Two tothree sets of electrochemical tests were conducted in eachexperimental condition, and all the OCP and potentiodynamicanodic polarization plots were highly reproducible.

The surface morphology and characterization of passivefilm compositions

The surface morphology of specimens after the corrosion testwas observed with a SEM, JEOL JSM-6510LA model, usingan acceleration voltage of 10–15 kV.

The XPS measurements were carried out on eachsample measured after 240 h immersion tests in 15.65 MHNO3. The XPS measurements were performed using aJEOL JPS-9200 model equipped with the dual X-raysource. The MgKα (1253.6 eV) X-ray was used togenerate the photoelectrons. The base pressure wasmaintained at 10−7 Pa. The X-ray gun was operated at100 W (10 kV, 10 mA) with a takeoff angle θ = 0o.Sputter depth profiles were evaluated using an argongun of Ar+ ion beam energy of 3 keV and a beam current

of 10 mA with sputter rate 2 nm/min determined from theSiO2 standard. The survey spectrum (0–1000 eV) wasrecorded for each sample to identify the elements present.The details of the XPS measurements and analysismethods have been described elsewhere [13, 14].

Results and discussion

The changes in the OCP of type 304L SS and type 310LSS measured at different concentrations of 1, 3, 6, 9, and11.5 M HNO3 are shown in Fig. 2a, b and Table 1. InFig. 2a, the OCP of type 304L SS shows active potentialin 1 M HNO3 (−0.024 ± 0.5 V vs. Ag/AgCl) and noblepotential for measurement in 3 M HNO3 (0.445 ± 0.1 Vvs. Ag/AgCl), 6 M HNO3 (0.670 ± 0.02 V vs. Ag/AgCl),9 M HNO3 (0.735 ± 0.01 V vs. Ag/AgCl), and 11.5 MHNO3 (0.815 ± 0.003 V vs. Ag/AgCl). Also, type 310LSS (Fig. 2b) shows the similar pattern with less noblepotential in 1 M HNO3 (0.078 ± 0.5 V vs. Ag/AgCl),followed by 0.590 ± 0.2 V vs. Ag/AgCl in 3 M HNO3,

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Fig. 2 Open circuit potentials vs. time behavior of alloy steels in differentconcentrations of 1, 3, 6, 9, and 11.5MHNO3: a type 304L stainless steeland b type 310L stainless steel

J Solid State Electrochem (2015) 19:3533–3542 3535

6 M HNO3 (0.710 ± 0.05 V vs. Ag/AgCl), 9 M HNO3

(0.750 ± 0.01 V vs. Ag/AgCl), and the most noble poten-tial in 11.5 M HNO3 (0.840 ± 0.02 V vs. Ag/AgCl). It canbe seen from these results that the OCP behavior isstrongly dependent on the nitric acid concentration. Incomparison, the OCPs of type 304L SS and type 310LSS are similar, except for 1 M HNO3 active potentialwhich was observed for type 304L SS. This indicates thatthe cathodic reactions occurring at the surface do not pro-vide the driving force and necessity for spontaneous pas-sivation, or implying that the thickness of all the filmswas being reduced [15, 16]. However, in both the alloysteels, the OCP is shifted to more noble potentials withincreasing nitric acid concentrations and similar behaviorshave been also previously observed [10, 16, 17]. Thechange in OCP to more noble potentials can be attributedto the faster growth rate of the passive film in this medi-um and thereby also in the thickness of the film [14–16].In addition, more noble potential with increasing nitricacid concentrations can also indicate strong polarizationand higher corrosion rate [13, 16]. The shift in OCP isalso affected by the change in both anodic and cathodicreaction rates. Thus, both decreased in the anodic reactionor increases in the cathodic reaction can lead to the shiftof OCP [18]. Furthermore, the OCP magnitudes can be agood macro surface indicator, and the results combinedwith more resolution surface techniques will be a key tothe investigation of nanometer to sub-nanometer compo-sitional fluctuation of passive films [19] at OCP to sup-port the above observation and mechanism. Based on ourpresent results, none of the investigated samples of types310L and 304L SS measured at OCP condition show nocorrosion product, and further studies in this area is es-sential to correlate the link between the corrosion prod-ucts formed vs. OCP. Furthermore, due to highly oxidiz-ing nature of nitric acid at higher nitric acid mediums, this

shift is undesirable, because of risk of overshooting thepotential beyond the passive region as these are associatedclose to the transpassive state potential [13, 15–17].

Potentiodynamic anodic polarization measurements

The potentiodynamic anodic polarization curves of type304L SS obtained for different concentrations of 1, 3, 6,9, and 11.5 M HNO3 are shown in Fig. 3a. The potenti-odynamic anodic polarization plot parameters such as cor-rosion potential (Ecorr), passive current density (ipass), andtranspassive/breakdown potential (ETP) are listed inTable 2. As observed in Fig. 3a, with the increase in nitricacid concentrations, (i) the corrosion potential, Ecorr, isshifted to more noble potential with increasing concentra-tions, (ii) the differences in ipass obtained at potential∼0.85 V vs. Ag/AgCl of the corresponding passive re-gions were not significant, and (iii) ETP marginally in-creases with concentrations. The gradual shifts in Ecorr

from −0.046 ± 0.1 V V vs. Ag/AgCl in 1 M HNO3 to0 . 390 ± 0 .05 V vs . Ag /AgC l i n 3 M HNO3 ,0 . 610 ± 0 .01 V vs . Ag /AgC l i n 6 M HNO3 ,0.706 ± 0.05 V vs. Ag/AgCl in 9 M HNO3, and0.765 ± 0.01 V vs. Ag/AgCl in 11.5 M HNO3 are ob-served. Furthermore, the transpassive/breakdown poten-tial, ETP terminating with a sharp rise in the current den-sity (Table 2), increases marginally with increasing nitricacid concentrations (i.e., 0.93 ± 0.1 V vs. Ag/AgCl for1 M HNO3, 0.99 ± 0.2 V vs. Ag/AgCl in 3 M HNO3,1.02 ± 0.05 V vs. Ag/AgCl in 6 M HNO3, 1.06 ± 0.1 Vvs. Ag/AgCl in 9 M HNO3, and 1.10 ± 0.1 V vs. Ag/AgClin 11.5 M HNO3). However, the measured ipass (Table 2)are not significantly affected by the nitric acid concentra-tions. Similarly, the potentiodaynamic anodic polarizationplots of type 310L SS shown in Fig. 3b also showed theshifts in Ecorr of 0.03 ± 0.1 V vs. Ag/AgCl in 1 M HNO3

to 0.480 ± 0.5 V vs . Ag/AgCl in 3 M HNO3,0 . 5 75 ± 0 . 1 V v s . Ag /AgC l i n 6 M HNO3 ,0.740 ± 0.1 V vs. Ag/AgCl in 9 M HNO3, and0.825 ± 0.1 V vs. Ag/AgCl in 11.5 M HNO3, respectively.Thereby, it is clear that the shifts of Ecorr to more noblepotentials with increasing concentrations are similar tothose of the shift in OCP (Fig. 2a,b) observed in boththe alloy steels. Furthermore, the ipass (Table 2) near to-wards ETP are low (∼106 A/cm2). Similarly, as observedin type 304L SS (Fig. 3a), the ETP also increases withconcentrations (Table 2). In both the alloy steels, the no-ble potential shifts of OCP (Fig. 2) and Ecorr (Fig. 3) withincreased in nitric acid concentrations may result intotranspassive corrosion. This is consistent with our previousresults [13–15] and reported by others [2, 3, 16, 17], and maybe explained as follows: As nitric acid is an oxidizing agent,the shift of Ecorr to more noble potentials are dependent on the

Table 1 The measuredOCP of type 304L andtype 310L stainless steelsmeasured at differentnitric acid concentrations

Nitric acid OCP(V vs. Ag/AgCl)

304L stainless steel

1 M HNO3 −0.024 ± 0.5

3 M HNO3 0.445 ± 0.1

6 M HNO3 0.670 ± 0.02

9 M HNO3 0.735 ± 0.01

11.5 M HNO3 0.815 ± 0.03


1 M HNO3 0.078 ± 0.5

3 M HNO3 0.590 ± 0.2

6 M HNO3 0.710 ± 0.05

9 M HNO3 0.750 ± 0.01

11.5 M HNO3 0.840 ± 0.02


autocatalytic reduction of nitric acid; thereby with increasingconcentrations, higher autocatalytic contributions are

expected via the global reduction of nitric acid [2, 13, 16].In the nitric acid, the cathodic reaction can be represented asfollows [3, 7, 20]:

HNO3 þ Hþ þ e−→NO2 þ H2O ð2Þ

Similarly, the reduction of nitrate which imposes the redoxpotential of the solutions in most of the nitric media can berepresented as follows [2, 16, 21]:

NO3− þ 3Hþ þ 2e−→HNO2 þ H2O ð3Þ

HNO2 þ Hþ þ e−→NOþ H2O ð4Þ

The dominant corrosion reactions responsible for corrosionof SS in nitric acid are attributed to the depolarizing cathodicreaction and the decomposition of unstable HNO2 that leads toa shift in Ecorr and transpassive corrosion [2, 16, 17]. The shiftin Ecorr to more noble potential is also expected as the redoxpotential of the medium increases with an increase in the con-centration of nitric acid [2, 13, 16, 17]. Furthermore, the ca-thodic reaction is largely nitrate reduction and generates oxi-dant (e.g., NO2, HNO2, NO, etc.) by the autocatalytic reduc-tion of nitric acid whereby the reduction rate increases withincreases in concentrations consequently leading to theshifting of the Ecorr of the steel to more noble potential [2, 3,15, 16]. The much larger Ecorr shift is also attributed to thenitrate reduction, leading to higher Ecorr [16]. However, de-pending on the concentrations of nitric acid, temperature, andions present, the end reaction products are variable [2, 3, 5].Therefore, increasing the acid concentrations accelerates nitricacid reduction and thereby redox potential increases [2, 7, 16].Additionally, in nitric acid corrosion, a large amount of HNO2

that is thermodynamically unstable, promotes the local attacksthat are formed by the electrochemical reaction in boilingHNO3. Thus, an increase in Cr content is required to obtain

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Fig. 3 Potentiodynamic anodic polarization curves of alloy steels at thescan rate of 0.167 mV s−1 in different concentrations of 1, 3, 6, 9, and11.5 M HNO3: a type 304L stainless steel and b type 310L stainless steel

Table 2 The potentiodynamicanodic polarization parametersobtained for type 304L and type310L stainless steels in differentnitric acid concentrations

Nitric acid Ecorr

(V vs. Ag/AgCl)ipass (A cm−2) ETP

(V vs. Ag/AgCl)


1 M HNO3 −0.046 ± 0.1 2.10 ± 0.03 × 10−6 0.93 ± 0.1

3 M HNO3 0.390 ± 0.05 1.88 ± 0.05 × 10−6 0.99 ± 0.2

6 M HNO3 0.610 ± 0.01 1.75 ± 0.02 × 10−6 1.02 ± 0.05

9 M HNO3 0.706 ± 0.05 1.72 ± 0.05 × 10−6 1.06 ± 0.1

11.5 M HNO3 0.765 ± 0.01 1.75 ± 0.03 × 10−6 1.10 ± 0.1


1 M HNO3 0.03 ± 0.1 1.16 ± 0.1 × 10−6 0.94 ± 0.1

3 M HNO3 0.480 ± 0.5 1.26 ± 0.05 × 10−6 0.98 ± 0.01

6 M HNO3 0.575 ± 0.1 1.02 ± 0.02 × 10−6 1.00 ± 0.01

9 M HNO3 0.740 ± 0.1 1.09 ± 0.02 × 10−6 1.06 ± 0.01

11.5 M HNO3 0.825 ± 0.1 8.60 ± 0.05 × 10−7 1.07 ± 0.01


stable oxide film and to eliminate the deleterious effect ofimpurity Si in the passive film region. The presence of Sipolarizes cathodic reactions that are the reduction of HNO3

and oxidizing ions like Cr(VI), resulting in the shifting of Ecorr[22]. Si is one of the most effective alloying element for im-provement corrosion resistance in highly oxidizing nitric acid:the presence of Si increases the corrosion resistance in highlyoxidizing nitric acid solutions containing Cr(VI) ions by sup-pressing cathodic reaction and by mitigating the detrimentaleffect of P segregated intergranularly [3, 11, 12, 22].

Corrosion mechanism of nitric acid concentrationson the passivity

The high corrosion resistance of SS in nitric acid is relatedto the stability of the passive films containing mainly chro-mium oxide/hydroxide [1–3, 23, 24]. For an in-depth un-derstanding of the nature of passivity, elucidation of themechanisms involved in the formation of passive films isof major significance. The passive film formation mecha-nism of stainless steel is a complex process, including sev-eral reaction stages: formation and growth of a passive filmvia generation, transport and annihilation of ionic pointdefects, continuous changes in the stoichiometry of thisfilm and especially of the first atomic layers adjacent tothe electrolyte , charge– t ransfer react ions at thefilm/electrolyte interface, and transport of reaction prod-ucts in the bulk solution [23–25]. Similarly, the stabilityof passive film (i.e., surface oxide) varied in a wide rangewith respect to the alloy composition, environment, filmthickness, structure, stoichiometry, electronic band struc-ture, ionic conductivity, temperature, passivation time,etc. [23, 24]. The influence of nitric acid concentrationon passivity behavior is best illustrated, by the schematiccurrent–voltage or polarization behavior shown in Fig. 4.As evident in Figs. 3 and 4, the gradual shift of Ecorr isrepresented by Ecorr1, Ecorr2, Ecorr3, Ecorr4, and Ecorr5,depicting thereby the influences of nitric acid concentra-tions on the passivity range and corrosion potential(Fig. 4). In lower concentrations (i.e., Ecorr1 and Ecorr2),both alloy steels (Fig. 4) are passivated over a wide poten-tial range and narrow region at higher concentrations, butits passive current is extremely low (Table 2). This behav-ior indicated that increases in concentrations have an effecton the corrosion resistance in both the alloy steels, as itresulted in an increase in Ecorr to more noble potential.Furthermore, at higher 11.5 M HNO3, the Ecorr5 (Fig. 4)is close to the transpassive state potential; these conditionsproduce changes in the composition of the passive layer,and thus, the passive layer is less stable [2, 3, 7]. The shiftof Ecorr with nitric acid concentrations is attributed to in-creases in the oxidizing power of HNO3 which may result

into transpassive corrosion. The transpassive corrosion(i.e., in the case of nitric acid corrosion at high concentra-tion) or transpassivation is a phenomenon observed in SSin which when it is exposed to highly oxidizing environ-ments (i.e., higher concentrations, temperatures, and oxi-dizing species) the corrosion potential is shifted to the no-ble direction (Fig. 4) near or beyond the passive region andmanifests itself as an intergranular corrosion, owing to thepreferential attack along grain boundaries. The shift ofEcorr to more noble potential is also attributed to strongpolarization with increasing nitric acid concentrationwhere passivity is lost and high corrosion rates can result[2, 7]. Furthermore, if the medium becomes excessivelyoxidizing, i.e., increase concentrations, passive film disso-lution occurs, especially by oxidation of Cr(III) (as insol-uble Cr2O3) to Cr(VI) (as soluble Cr2O7

2−) [1, 11, 16, 24].The transfer of oxidized ions into the electrolyte also de-polarizes the cathodic reaction as it is a strong oxidizer,contributing to a more severe attack and autocatalytic pro-cess conditions (Reactions 2–4) [17]. This results in accel-erating transpassive corrosion.

There exist many models that provide an experimentaland theoretical basis for describing the transpassive mech-anism for passive film breakdown/dissolution. The energyband model proposed by Sato [26] correlated thetranspassive dissolution kinetic to the relative positionsof the Fermi level and the conduction band in the passivefilm [26]. The point defect model (PDM) by Macdonald[27] assumed that the passive film is a highly defectivebarrier film with point defects resulting from metal andoxygen vacancies and metal interstitial in which the va-cancies act as the electronic dopants. The vacancy gener-ated and annihilated by reactions at the interface leads tofilm decohesion and breakdown. According to Bojinov

Fig. 4 Schematic illustration of anodic polarization curve depicting theinfluences of HNO3 concentrations on the passivity range and corrosionpotential near transpassive potential


et al. [25, 28], transpassive dissolution in SS occurs viaCr(VI) intermediate species. In Ni-based alloys, Ni disso-lution occurs via two parallel reaction paths, featuring twoadsorbed intermediates. Fe acts as a secondary passivatingagent and Mo as an accelerator of transpassive dissolution[28]. It also employed a curve-fitting technique to explainthe transpassive dissolution using the mathematical meth-od and electrochemical parameters. Song [29] proposed adistorted non-stoichiometric oxide film model for Fe-Cr-Ni stainless steels where the transport and migration ofions in the passive film is not limited by the vacancy,but transpassivation is associated closely with the disso-lution of Cr and Ni of the film. The model also explainsthe changes in film composition and film thickness affect-ing the transpassive dissolution. In our present work,nickel was not found in the passive film or as a corrosionproduct. In addition, the role of Ni affecting passivity/transpassivity continues to be a matter of debate [23, 24,30, 31]. Vignal et.al [30] detected nickel at the substrate/passive film interface in duplex stainless steels and alsoattributed that nickel oxide can be formed at high anodicpotentials and low pH values during electrochemical etch-ing in nitric acid. In AISI types 430, AISI 304L and AISI316 SS, Ni in the alloy is known to increase the Cr/Feratio and affects Cr distribution across the passive films,determining the film growth and the corrosion resistance[31]. In the transpassive range, a transition from Ni(II) tosoluble Ni(III) occurs and the layer dissolves. However,the Ni(II) oxide and hydroxide in the passive oxide filmprotect the substrate metal against generalized corrosion[32]. The presence of increasing Ni content in Fe-Cr-Nisteels is also known to promote the formation of a thinnerand more protecting passive film. Ni(II) in the passivefilm catalyzes O2 and H2 evolution [31], which subse-quently can change the cathodic reaction, with respect toEcorr, but this reaction can also be influenced by the het-erogeneous distribution of the corrosion products at thesurface.

The control step mechanism at the interface based on thenitric acid and its influence on the transpassivation dependedon many factors, i.e., alloy composition, solute segregation,impurities, heat treatment, non-metallic inclusion, microstruc-ture, etc. [1–3, 22, 33]. Extensive studies also show that theIGC in nitric acid is related to the structure and energy of thegrain boundaries that result to the formation of active sites [34,35]. The grain boundary engineering (GBE) using the coinci-dence site lattice (CSL) concept to prevent the initiation andpropagation of intergranular degradation at the interface andalong the grain boundaries was also recently highlighted[35, 36]. A twin-induced GBE with the high CSL frequencyin type 304 SS showed a much higher resistance totranspassive intergranular corrosion during the Coriou testattributed to a very low percolation probability of random

boundary networks in the over-threshold region andremarkable suppression of intergranular deterioration [36].Hence, many factors influence the control step mechanismat the interface and thereby affecting the transpassivationbehavior in alloy steels.

Corrosion rate in boiling nitric acid

The corrosion rate of type 304L SS measured in boiling15.65 M HNO3 for five 48 h cycles for a total period of240 h by the weight loss method is shown in Fig. 5. Thecorrosion rate increases with an increased in immersion time.The corros ion ra te shows an ini t ia l increase of0.18 ± 0.02 mm/y in 48 h, 0.19 ± 0.001 mm/y in 96 h, andthen 0.2 ± 0.01 mm/y after 192 to 240 h. However, the mea-sured corrosion rates of the type 310 L SS (Fig. 5) showedalmost a cons tan t and low corros ion ra te ( i .e . ,0.006 ± 0.001 mm/y up to 240 h). In austenitic SS, the corro-sion resistance in highly oxidizing nitric acid solution isstrongly affected by Cr content and impurity of the residualelements segregated at the grain boundary [1–3, 24]. Thereby,these results indicated that type 310L SS with higher Cr(25 wt%) and Ni (20 wt%) and lower Si content (0.2 wt%)results in improvement in the corrosion resistance in nitricacid. Furthermore, the differences on the corrosion rate be-tween the two alloy steels can also be corroborated by com-paring the SEM morphology of the specimens after the corro-sion test (Fig. 6a,b). The surface morphology (Fig. 6a) afterthe boiling test for 240 h of type 304L SS shows severe IGCattack, i.e., a preferential attack along grain boundaries. In thetype 310L SS, intensity of IGC attack (Fig. 6b) is much lower.In nitric acid, the IGC attack in SS results from the segregationof impurities (i.e., C, P, S, B, Si, Mo, N, etc.) in the γmatrix atgrain boundaries. Such agglomeration may result in

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Fig. 5 The measured corrosion rate of type 304L stainless steel and type310L stainless steel in immersion/boiling test in 15.65 M HNO3 after240 h


precipitation of chromium carbides, with residual elementslike P, N, Si, or S replacing Cr or C, and the resultant forma-tion of active sites on the surface [12, 33]. Also, whethercorrosion is predominantly by IGC attack or general corrosionattack depends upon the difference to the rate of corrosion ofthe grain boundary zones and the grain faces [3, 8, 34]. Theincrease in corrosion rate with increasing immersion time intype 304L SS may also be attributed to lower Cr content andthe segregation of the above mentioned impurity/alloying el-ements in grain boundaries that enhanced IGC attack [3]. Inaustenitic SS, higher chromium along with Ni of more than20 % is also known to improve corrosion resistance in nitricacid with oxidizing ions [3, 22], thereby inhibiting the IGCgrain boundaries’ attack. It is essential to adjust major ele-ments of Cr and Ni to be upper/higher level and control Si0.1–0.2 wt%, limit residual/minor element (P, S, B, N, etc.),and modify the microstructure [3, 12]. Hence, high Cr and Niin type 310L SS are designed for improving the transpassive

corrosion resistance and optimizing elements of Cr, W, and Si,which are oxide film formers and cathodic depolarizers [1, 2,22]. As observed in the present results, better corrosion resis-tance clearly seen in type 310L SS may provide good reliabil-ity of SS used in reprocessing plant.

Passive oxide film compositions

The XPS surveyed spectra of the passive oxide film formed ontype 304L SS and type 310L SS after 240 h of immersion testsare shown in Fig. 7. In Fig. 7, the spectra of type 304L SSshows the presence of mainly chromium (Cr 2p), iron (Fe 2p),oxygen (O 1s), silicon (Si 2s and 2p), and carbon (C 1s)signals. The wide surveyed spectra of type 310L SS (Fig. 7)are similar and also show the chromium (Cr 2p), iron (Fe 2p),oxygen (O 1s), silicon (Si 2s and 2p), and carbon (C 1s)signals. Both the alloy steels do not show Ni peak intensity.However, the intense peak intensities of chromium (Cr 1p and2p), oxygen (O 1s), and silicon (Si 2s and 2p) in both alloysteels (Fig. 5) may indicate its enrichment discussed later. Inboth the steels, the chromium peak (Cr 2p) contains at leasttwo contributions, including a large oxidized signal at∼576 eV corresponding to Cr(III) of Cr(III) oxides and hy-droxides (∼577 eV). The iron peaks (Fe 2p) at binding energyat ∼710 and ∼712 eV, corresponding to Fe(III) oxide (Fe2O3)and Fe(III) hydroxide (Fe(OH)3). The O1s peak indicates thatoxygen is present under different chemical states such as O2−

(∼530 eV), OH− (∼531 eV) and adsorbedH2O (∼533 eV). Thenon-deconvoluted Si 2p (102.6 eV) and Si 2s (150.51 eV)peaks are attributed to the SiO2 [11, 23, 24]. The C1s signalat ∼284 eVis characteristic of the contaminants on the surface.The determination of the detail oxidation states of the ele-ments composing the passive films is a complex operationthat goes beyond the scope of this paper.

Fig. 6 Typical SEM morphology showing intergranular corrosion attackin alloy steels after boiling test in 15.65 M HNO3 of 240 h exposure: atype 304L stainless steel and b type 310L stainless steel

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Fig. 7 Typical surveyed spectrum of XPS measurements obtained foralloy steels of type 304L and type 310L stainless steels after boiling testin 15.65 M HNO3 of 240 h exposure


The XPS sputter depth profiles of Cr 2p, O 1s, Fe 2p,C 1s, Si 2s, and 2p obtained for the type 304L SS and310L SS after a 240-h boiling test are shown in Fig. 8a,b,respectively. In (Fig. 8a), the depth profiles of type 304LSS show that the atomic percent of Cr, C, and Si contentdecreases during the initial stage of the sputtering periodand maintained a constant value, and C is present, mainlyas a surface contaminant. On the other hand, the atomicpercent of Cr, Si, and O increases with sputtering timeand remains almost a constant value with sputtering time.The increased concentration of O indicated the presenceof oxygen as a typical bulk in the specimens.

In type 310L SS, (Fig. 8b) the sputter depth profiles of Cr2p, O 1s, Fe 2p, C 1s, Si 2s, and 2p showed similar to thepattern observed in type 304L SS (Fig. 8a). However, low Feconcentration (5 at.%) in both alloy steels may be attributabledue to higher dissolution. Similarly, the Cr atomic percentconcentration differences between the two alloy steels arealmost in a similar range. Marginally lower Cr concentrationin the passive film of both the steels may be attributed to Cr

predominating the competitive oxidation and thereby diffusesto a greater extent into the film [13, 23]. The enrichments of Siin a range ∼10–15 at.% (Si 2s) and ∼5–6 at.% (Si 1s) areclearly evident, thereby indicating the presence of oxidizedelements. As described earlier, silicon is considered as oneof the most important alloying elements in austenitic SS toimprove their corrosion resistance in strongly oxidizing nitricacid environments and to inhibit the IGC [11]. However, therole of silicon is less obvious, and literature is still not clear onthe beneficial role of Si improving the corrosion resistance andthe passive layer property [1, 5, 11, 33, 34]. Some authorsattributed to enrichments of SiO2 that suppresses the cathodicreaction and, consequently, decreases the corrosion rate [11,12]; silicon in the passive film decreased the chromium con-tent of the passive layer [11], others suggest the presence in theform of silicate [11] or dual layer film (i.e., outer Cr2O3 andinner SiO2) [37]. However, present results confirm the pres-ence of SiO2 in the passive layer that may further enhance theprotective ability of the Cr2O3 [11, 12]. However, the role ofSi and its influence on corrosion resistance in nitric acid needmore elaborate and further investigation.

Conclusions

The corrosion resistance of type 304L SS and type 310L SSwas studied in different concentrations of 1, 3, 6, 9, and11.5 M HNO3, and boiling 15.65 M HNO3 for 240 h. Theelectrochemical corrosion test results showed similar behav-iors in both alloy steels: the open circuit potentials along withcorrosion potential are shifted to more noble potential withincreasing concentrations from 1 to 11.5 M HNO3.However, the passive current density was not significantlyaffected, but the transpassive potential increases marginallyto a higher potential with concentrations. The shift in OCPand corrosion potential to more noble potentials can be attrib-uted to the faster growth rate of the passive film in this medi-um. The corrosion rate measured in boiling 15.65 M HNO3

for 240 h show much lower corrosion rate in type 304L SS(∼0.06 ± 0.012 mm/y) then type 304L SS (∼0.18 ± 0.02–0.2 ± 0.001 mm/y). The SEMmorphologies corroborated the-se results that show severe IGC in type 304L SS than type310L SS after 240 h of exposure. The XPS analysis of thepassive oxide films of both alloy steels appears similar thatshows the presence of Cr2O3 and SiO2. Present results indi-cated that the corrosion resistance of type 310L SS was supe-rior to type 304L SS in boiling nitric acid. The improvedcorrosion resistance of type 310L SS was attributed to controlchemical composition in the alloy and enrichment of Si layerin addition to Cr2O3 in the passive film. Thereby, better cor-rosion resistance of type 310L SSmay provide good reliabilityfor application in the reprocessing plant.

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Fig. 8 XPS depth profile alloy steels after boiling test in 15.65 M HNO3

of 240 h exposure: a type 304L stainless steel and b type 310L stainlesssteel


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