The Inhibition of Hydrogen Embrittlement in SAE 4340 Steelin an Aqueous Environment with the Rare Earth CompoundLanthanum 4 Hydroxy Cinnamate

B.R.W. HINTON, M. BEHROUZVAZIRI, M. FORSYTH, R.K. GUPTA, M. SETER,and P.G. BUSHELL

A preliminary study showed that the inhibitor lanthanum 4-hydroxy cinnamate ((La4OHcin)3)at a concentration of 400 ppm prevented the hydrogen embrittlement (HE) of SAE 4340 steeltensile specimens when tested under slow strain rate conditions in a 0.01M NaCl. In the presenceof the inhibitor, a complex film formed on the surface of specimens during the slow strain ratetest (SSRT), and no corrosion pits were detected. Electrochemical polarization studies indicatedthat the La(4OHcin)3 acted as an anodic inhibitor in the NaCl solution. This article alsodiscusses the mechanism of HE inhibition by La(4OHcin)3.

DOI: 10.1007/s11661-012-1103-y� The Minerals, Metals & Materials Society and ASM International 2012

I. INTRODUCTION

CHROMATES have long been considered one of themost effective corrosion inhibitors. However, because ofsignificant environmental and health concerns,[1] there hasbeen considerable focus on nontoxic alternatives such asrare earth metal compounds.[2–4] In 1988, Hinton et al.[5]

showed that the addition of 100 ppmof cerium chloride toflowing tap water reduced the corrosion rate on mild steelby a factor of 10. In unrelated work, the inhibitingproperties of organic salts were also investigated byMercer,[6] who demonstrated that the alkali metal saltsof carboxylic acids such as cinnamates effectively inhibitedthe corrosion of mild steel. Forsyth and co-workersconsidered these unrelated findings and synthesized com-pounds consisting of an organic anion (e.g., carboxylateion) and a rare earth metal cation that could produce asynergistic inhibiting effect.[7–13] In particular, Blinet al.[7,10] and Behrouzvaziri et al.[13,14] showed usingelectrochemical studies that lanthanum 4-hydroxy cinna-mate (La(4OHcin)3), when present at levels up to 900 ppmin 0.01M NaCl, significantly increased the polarizationresistance ofmild steelwith exposure time compared to thecontrol condition. The inhibitor also prevented rustingand pitting during a 7 day immersion test.

Previous studies by Gerrard et al.[15,16] with the high-strength low-alloy tempered martensitic SAE 4340 steel

showed that this steel is very susceptible to corrosionwhen immersed in NaCl solutions at pH values between4 and 9, with pits developing in conjunction with a red-orange rust film.It has long been recognized that in aqueous environ-

ments under freely corroding conditions, such steels aresusceptible to hydrogen embrittlement (HE).[17–19] Therole of the environment is to allow cathodic reactionssuch as the reduction of hydrogen (H) ions or thereduction of water to occur, which results in thegeneration of H atoms at the steel surface and theirsubsequent diffusion into the steel.[20] Once in the steel,under load, H will diffuse to local stress raisers such assurface irregularities (scratches, notches), changes ofsection, corrosion damage including pits, and certainmicrostructural features. The mechanisms throughwhich the hydrogen degrades the steel and reduces itsload carrying capacity are still open to debate and arecontroversial, even after decades of research. They haverecently been reviewed by Gangloff et al.[17,18]

Given the effectiveness of La(4OHcin)3 as a corrosioninhibitor for mild steel,[8,10] it was of interest toinvestigate if it could inhibit the corrosion of SAE4340 steel and thereby prevent HE. This article reportson an initial study to assess the effectiveness ofLa(4OHcin)3 in preventing the HE of SAE 4340 steelwhen tested under slow strain rate conditions in a NaClsolution.

II. EXPERIMENTAL DETAILS

A. Steel and Specimens

For this study, the notched tensile specimens(Figures 1(a) and (b)) were manufactured from a singlebatch of SAE 4340 steel bars used not only for this workbut also for previous studies.[21] The bars had acomposition (wt pct) of 0.43C, 0.84Mn, 0.29Si, 1.78Ni,

B.R.W. HINTON, Adjunct Professor, is with the Department ofMaterials Engineering, Monash University, Wellington Road, Clay-ton, VIC 3800, Australia, and is also with the Institute for Technology,Research and Innovation, Deakin University, Burwood Highway,Burwood, VIC 3125, Australia. Contact e-mail: bruce.hinton@monash.edu M. BEHROUZVAZIRI, Student, and R.K. GUPTA, ResearchFellow, are with the Department of Materials Engineering, MonashUniversity. M. FORSYTH, Professor, and M. SETER, ResearchFellow, are with Institute for Technology, Research and Innovation,Deakin University. P.G. BUSHELL, Scientist, is with the Departmentof Defence, Defence Science and Technology Organisation, MaritimePlatforms Division, Lorimer St., Fishermans Bend, VIV 3207, Australia.

Manuscript submitted May 9, 2011.Article published online March 27, 2012

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 43A, JULY 2012—2251

0.85Cr, 0.25Mo, 0.009S, <0.01P, <0.01Al, and balanceFe. Heat treatment involved austenitization (molten saltbath) at 1088 K (815 �C) for 1 hour followed byquenching in oil at 323 K (50 �C) and tempering(molten salt bath) twice at 533 K (260 �C) for 1 hourwith air cooling between tempers. This produced atempered martensitic microstructure. The notch wasprepared by low stress crush grinding (at a maximum of423 K (150 �C)) after heat treatment and was then stressrelieved with a bake at 463 K (190 �C) for 3 hours in aninert environment. The ultimate tensile strength of thesteel after this heat treatment was 1845 MPa.[21] TheVickers hardness was measured at seven points on a2000 mesh abraded steel surface with a 2 kg load. Thevalue was 573 with a 1SD of ±9.

Figure 2 is a polished and etched (2 pct Nital 10 sec-onds) metallographic section taken from a plane normalto the axis of a tensile specimen. The microstructureconsisted of fine tempered martensite typical of this steelheat treated at such high-strength levels. The prioraustenite grain size was determined in accordance withASTM E112-10 Standard Test Methods for Determin-ing Average Grain Size, after etching the surface with apicric acid-copper chloride solution. The average prioraustenite grain size was 12 ± 1 lm.[15]

B. Slow Strain Rate Tensile Tests

The slow strain rate test (SSRT) or constant exten-sion rate test used in this study, where a tensilespecimen is loaded in the presence of the environment,is a very effective tool for quantifying the degree of HEinduced in high-strength steels and other alloys by thatenvironment.[21–24] The principal advantage of this testis that the susceptibility to HE of a particular steel–environment combination is rapidly assessed, andspecific parameters that quantify the HE effect areobtained.[25] These parameters include time to failure,reduction in fracture stress, and reduction in ductil-ity.[25] The accepted standard for HE testing ASTM F519[26] is based on the use of notched tensile specimensmade from AISI 4340 air melted and heat treated to1790 to 2070 MPa ultimate tensile strength. In thatcondition, the steel was shown to be most susceptibleto HE. The notch concentrates stress, limits ductility,and localizes fracture and any effects of hydrogenuptake.

SSRTs were conducted at a constant crossheaddisplacement rate of 2 9 10�6 mm s�1 on a computer-controlled tensile test machine described elsewhere.[21] Ithas been shown by Pollock[21,27] that at this displace-ment rate, the embrittling effect of an aqueous environ-ment on SAE 4340 steel is the most pronounced,although more recent work by Ramamurthy et al.showed that the degree of HE is even greater with lowerstrain rates.[23,24] A strain rate of 2 9 10�6 s�1 wassuggested as suitable[25,28] for assessing the ability of anenvironment to induce HE. The effects of displacementrate in the presence of La(4OHcin)3 inhibitor will be thesubject of a future publication. Prior to testing, thespecimens were rinsed with ethanol and dried in air. Allbut the gage length of the specimens was coated with aprotective lacquer LACOMIT.* They were sealed into a

Teflon cell (Figure 3), which allowed the ends of thespecimen to be attached to the loading grips on thetensile test machine. The cell contained 50 mL of testsolution open to air for the test duration, and all testswere conducted at the open circuit potential (OCP). Theapplied load at the point of fracture was recorded usinga load cell. The fracture stress was calculated using the

Fig. 1—(a) Drawing and (b) photograph of the tensile test specimen. Dimensions are in millimeters.

Fig. 2—Polished and etched metallographic section of heat-treatedSAE 4340 steel (etchant: 2 pct Nital, 10 s).

*LACOMIT is a trademark of W. Canning & Co Ltd, Sheffield,UK.

2252—VOLUME 43A, JULY 2012 METALLURGICAL AND MATERIALS TRANSACTIONS A

load at fracture and the original cross-sectional area atthe root of the notch.

C. Electrochemical Polarization Tests

The working electrodes used for electrochemicalpolarization studies were machined from areas ofbroken tensile specimens outside the gage length. Theywere encapsulated and sealed in Teflon tubing, with awire attached for connection to a potentiostat. Thecross-sectional area of the tensile specimen formed theworking electrode surface (7.3 mm diameter and42 mm2 area). The electrode surface was abraded to a1200 mesh finish using silicon carbide papers. Thepolarization experiments were carried out in a standardthree electrode glass cell[29] with 250 mL of test solution.Potentials were measured relative to a saturated calomelelectrode (SCE) via a Luggin probe. A high surface area0.9 mm titanium mesh was used as a counter electrode.All electrochemical tests were conducted using a com-puter-controlled multichannel VMP3 Potentiostat andEC-Lab� Software (BioLogic, Claix, France). Theworking electrode was immersed in the test solutionfor 30 minutes before polarization was commenced.Some tests were started from potentials more negativethan the OCP, while others were started at the OCP. Thepotential was scanned at a rate of 0.167 mV/s. Freshlyprepared specimens were used for each scan.

The corrosion current (Icorr) values were calculatedusing well-known Tafel extrapolation methods. Thenear-linear and symmetrical sections of the anodic andcathodic curves within ±20 mV from Ecorr (indicatingcorrosion under activation control) were extrapolateduntil the lines intersected at Ecorr to give the Icorr value.

D. Test Solutions

The 0.01M NaCl control test solution was preparedusing analytical reagent grade sodium chloride anddistilled water. The solutions of La(4OHcin)3 wereprepared following the method outlined by Deaconet al.[30] Aqueous solutions (standardized by EDTAtitrations) of LaCl3Æ7H2O were prepared by dissolving

the LaCl3Æ7H2O solid in distilled water. An aqueoussolution of sodium 4-OH cinnamate (three equivalents)was slowly added to the LaCl3Æ7H2O solution, and aprecipitate formed instantly. Upon addition of sodiumsalt, the solution was adjusted to pH 5, stirred for1 hour, and filtered, and the precipitate was washed withdistilled water and dried in a vacuum desiccator for2 days. A schematic of the 4-hydroxy cinnamate anion isshown in Figure 4. In order to help dissolution of theinhibitor, 10 to 20 mL of 96 pct analytical grade ethanolwas added to the 0.1M NaCl solution. The ethanol wasthen evaporated by heating the solution at 353 K(80 �C) for an hour. All the solutions were left unstirredand open to air, in a temperature-controlled laboratoryatmosphere at 296 K (23 �C).

E. Microscopy

Fractography and scanning electron microscopy(SEM) were carried out on a JEOL** JSM-6490LA

scanning electron microscope. An Olympus model MX50 metallurgical microscope was used for the opticalmicroscopy.

III. RESULTS

A. Slow Strain Rate Tensile Tests

Figure 5 shows typical stress vs time data for steelspecimens loaded in 0.01M NaCl and 0.01M NaCl with400 ppm La(4OHcin)3. The data show that very littleplastic deformation occurred with the notched tensilespecimen before fracture. The nonlinearity near the start

Fig. 3—White Teflon test cell positioned in the slow strain rate ten-sile test machine.

Fig. 4—The 4 Hydroxy cinnamate anion.

0

500

1000

1500

2000

2500

3000

0 60 120 180 240 300

Time hours

Str

ess

MP

a

0.01M NaCl 0.01M NaCl + 400ppm Inhibitor

Fig. 5—Stress vs time data for specimens tested in a NaCl solutionwith and without inhibitor.

**JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 43A, JULY 2012—2253

of each test is associated with some ‘‘softness’’ of thegear train in the tensile machine under initial loading.

Figure 6 shows the fracture stress for specimensloaded in 0.01M NaCl, and 0.01M NaCl with 200 and400 ppm La(4OHcin)3. The individual points are theaverage of two results, and the error bars represent thesetwo results for each concentration. The results for thetests in air were 2410 and 2660 MPa. The valuesobtained in air are consistent with results obtained inprevious studies by Pollock and Hinton[31] and arewithin their scatter band obtained on 21 specimens.

The data in Figures 5 and 6 clearly show that whentested in 0.01M NaCl solution, the fracture stress for thesteel was reduced by more than 50 pct compared to thevalues obtained in air. With La(4OHcin)3 present insolution, the fracture stress increased with increasinginhibitor concentration. The fracture stress at 400 ppmof La(4OHcin)3 was similar to that obtained in air. Thetime to failure (Figure 5) in 0.01M NaCl was around80 hours, whereas with La(4OHcin)3 present, the time tofailure was around 210 hours. These data in Figures 5and 6 indicate that the presence of La(4OHcin)3 at400 ppm prevented the premature fracture of the steel.

It is noted that the values of the fracture stressobtained in air, i.e., 2410 and 2660 MPa, are signifi-cantly different from the value of the UTS of 1845 MPa.The UTS was obtained with smooth round tensilespecimens and not with notched round tensile speci-mens. The presence of the notch in a high-strength steelsuch as SAE 4340 introduces a triaxial stress state at thenotch tip,[32] and for the geometry of the specimens usedin our article, approximately plane strain conditionsapply. This stress state may cause an increase in the yieldstress at the notch tip up to values 3 times that of theuniaxial yield stress,[33,34] and an accompanying decreasein ductility in the zone ahead of the notch. This changein material properties just ahead of the notch tip is aresult of the effects of the triaxial stress state (i.e.,constraint arising from the notch geometry) on disloca-tion dynamics at the tip.[35] Dislocation motion requiresthe presence of shear stress. The triaxial stress statecannot produce a shear stress. As a result, dislocationmovement is restricted; hence, the yield strength and theultimate or fracture stress are increased (i.e., from 1845to 2408–2663 MPa). This restriction in dislocation

movement will cause the ductility to be reduced in thenotch region, and hence the absence of any evidence ofductility in the stress-time curves (Figure 5).

B. Fractography

The SEM of the fracture surface of specimens testedin air showed that microvoid coalescence (MVC) (dim-ples) and some flat transgranular areas of quasi-cleavagewere the dominant fracture modes (Figure 7) over mostof the fracture surface. MVC was the dominant mode onthe small shear lips around the edges of the fractureadjacent to the notch. These modes are typical offracture in air in this ultra-high-strength steel.[15] Thesurface of specimens tested in 0.01M NaCl had extensiveregions of intergranular fracture adjacent to the notch(Figure 8). These regions covered approximately 50 pctof the fracture surface with MVC and quasi-cleavagemodes over the remainder. There was also evidence ofductile tearing interspersed among and on the inter-granular facets. Intergranular fracture along prioraustenite grain boundaries is generally the dominant

500

1000

1500

2000

2500

3000

0 100 200 300 400 500

Inhibitor Concentration ppm

Fra

ctu

re S

tres

s M

Pa

Fig. 6—Effects of inhibitor concentration on fracture stress. Theindividual points are the average of two results, and the error barsrepresent the duplicate results for each concentration. The results fortests in air are 2408 and 2663 MPa.

Fig. 7—Scanning electron micrograph from the surface of a speci-men tested in air.

Fig. 8—Scanning electron micrograph from the surface of a speci-men tested in 0.01M NaCl.

2254—VOLUME 43A, JULY 2012 METALLURGICAL AND MATERIALS TRANSACTIONS A

mode of failure induced by HE in high-strength steelssuch as SAE 4340.[17] Intergranular fracture modes werenot detected on specimens fractured in air. Corrosiondamage, particularly pitting, was detected in the notchadjacent to the fracture surface (Figure 9). Many ofthese pits were associated with surface cracks below theplane of the fracture surface. The fracture surfaces ofspecimens tested in NaCl with 400 ppm La(4OHcin)3were almost identical to those tested in air with MVCand quasi-cleavage modes present and MVC around theedges of the notch circumference. However, some smallisolated pockets of intergranular fracture dispersedamong the areas of MVC were observed at a fewlocations around the notch circumference adjacent tothe notch tip (Figure 10). This observation suggests

some very localized HE may have occurred, butinsufficient to affect the fracture stress.

C. Optical Microscopy and SEM

At the conclusion of the SSRTs, the gage lengths andnotches of specimens were examined under an opticalmicroscope and in the scanning electron microscope.With specimens tested in NaCl, these areas were coveredwith a thin rust colored film, and several deep corrosionpits were observed (Figure 11). Energy dispersive X-rayspectroscopy (EDS) analyses (Table I) indicated thatthis film was predominantly Fe oxide. Some Cr andchloride ion was also detected in this film. In contrast,visual inspection of specimens tested in NaCl with theinhibitor present showed that the gage lengths andnotches were not rusted, but they did have a dullappearance. Under the high-magnification opticalmicroscope, it was clear that the dull appearance wasthe result of a film on the surface. This film had acomplex structure and consisted of many particles,which appeared either to be on top of, or emergentfrom, a background film, which covered the surface. Itwas very different from the rust films observed onspecimens tested in NaCl solution with no inhibitor.In the SEM, these particles were easily recognized(Figure 12), and several large particles had a definitecrystallographic appearance (Figure 13) similar to thoseobserved by Hinton et al. with mild steel tested in NaClsolution with CeCl3 present.

[5] One of these particles waspartially dislodged in handling prior to observation inthe SEM (insert Figure 12). The apparently pristinecondition of the abraded steel surface where the particlehad been can be seen. The fine nodular structure of thebackground film formed when the inhibitor was presentin the NaCl solution is shown in Figure 14. The cracksare the result of the film drying in the SEM. Nocorrosion pits could be detected within the resolutionlimits of the optical microscope or with the SEM on thespecimens tested in the presence of the inhibitor.Table I shows the results of EDS analyses of various

features of the film formed on specimens tested with400 ppm of the inhibitor present. These data show thatthe large particles (Figure 13), the general backgroundfilm (Figure 14(a), spectrum 1), and the smaller particles(Figure 14(b), spectrum 2) all contained Fe, La, and Owith C. The higher concentrations of La were in thelarger particles. These data strongly indicate that thefilm consists of a mixture of La and Fe oxide orhydroxide, with C indicating the probable presence ofthe cinnamate moiety. The reduced levels of Fe and Crin the film formed in the presence of the inhibitorcompared with the film formed with no inhibitor couldbe an indicator of a reduced amount of corrosionoccurring when the inhibitor was present.

D. Electrochemical Polarization Tests

Typical polarization curves obtained in solutionsopen to air are shown in Figure 15. The scans werecommenced at the OCP after 30 minutes in solution.Without inhibitor present, the OCP was around –0.5 V.

Fig. 9—Scanning electron micrograph of the notch of a specimentested in 0.01M NaCl showing corrosion damage and pitting(arrows) in the notch.

Fig. 10—Scanning electron micrograph from the surface of a speci-men tested in 0.01M NaCl with inhibitor present. The insert showsan isolated pocket of intergranular fracture.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 43A, JULY 2012—2255

With increasing inhibitor concentration, the OCP wasshifted in the positive direction. The cathodic curvesobtained with and without the inhibitor present indicatethat oxygen reduction was the dominant cathodicreaction, with the current density becoming almost

independent of potential. The effect of the inhibitor wasto shift the curve to higher potentials. Allowing for thehigher OCP when the inhibitor was present, these datashow that there was no effect of inhibitor on thecathodic kinetics. With inhibitor present, a clearlydefined pitting potential (Epit) was observed at around

Fig. 11—Scanning electron micrographs of the surface of a tensile specimen tested in 0.01M NaCl at 2 9 10�6 mm/s showing (a) corrosion pitsand (b) cracked corrosion product film. The cracks are the result of the corrosion product drying either in air or in the vacuum of the SEM.

Table I. EDS Data for the Films Formed on the Surfaces of (a) Specimen Tested in 0.01M NaCl at 2 3 1026 mm/s (No Inhibi-

tor) and (b) Specimen Tested in 0.01M NaCl with 400 ppm Inhibitor at 2 3 1026 mm/s (Inhibitor/Various Locations)

Location/Element (Wt Pct) Fe La Cr O C Cl

No inhibitor 83.1 — 3.1 13.9 — 2.1Inhibitor/large particle 31.6 25.3 0.1 25.2 12.9 0.9Inhibitor/background film 59.1 12.5 0.7 17.6 8.6 —Inhibitor/small particle 30.1 24.2 — 32.9 12.5 —

Fig. 12—Scanning electron micrographs of the surface of a tensilespecimen tested in 0.01M NaCl with 400 ppm inhibitor at2 9 10�6 mm/s showing a complex surface film consisting of manyparticles. The insert shows a large particle and a smaller particle,which has been dislodged.

Fig. 13—Scanning electron micrographs of the surface of a tensilespecimen tested in 0.01M NaCl with 400 ppm inhibitor at2 9 10�6 mm/s showing the crystalline nature of a large particle.

2256—VOLUME 43A, JULY 2012 METALLURGICAL AND MATERIALS TRANSACTIONS A

–0.2V for 200 ppm and around –0.1 V for 400 ppm. Theanodic current densities in the vicinity of the OCP werereduced by several orders of magnitude with inhibitorpresent. At very high potentials, the data suggest thatthe inhibitor had no effect on the anodic kinetics. Nowell-defined pitting potential was observed in theabsence of the inhibitor. In NaCl solution only, thissteel corrodes and undergoes pitting in a chloridesolution open to air,[15,16] with the OCP above or closeto the pitting potential. Typical values for OCP, Epit,and corrosion current (Icorr) obtained by Tafel extrap-olation are shown in Table II. Icorr data were obtainedfrom three polarization plots in many cases, and theresults in the Table II are shown as an average with ±1standard deviation. There was considerable scatter;however, significant trends are evident. With the inhib-itor present, the corrosion current was reduced by up to

an order of magnitude. Typical results from polarizationtests in 0.01M NaCl deaerated by bubbling nitrogen for30 minutes before testing are shown in Figure 16. Withthese tests, the scans were commenced at potentialsmore negative than the OCP. As in the tests with thesolution open to air, the presence of the inhibitor shiftedthe OCP to more positive potentials and had no effecton the kinetics of the cathodic reaction, which ishydrogen evolution. With the inhibitor present, the Epit

was increased and the anodic kinetics were reduced(Table II). The corrosion current in the absence of airwas also reduced (Table II) but not to the same degreeas in the presence of air. The lower corrosion currents inthe absence of air are consistent with the less aggressiveenvironment when oxygen is not present.The increase in Epit and the decrease in current density

in the pseudo-passive region of the anodic polarizationcurves between Ecorr and Epit with increasing La(4OH-cin)3 concentration is consistent with a thicker layer ofthe protective film forming on the surface at 400 ppm. Itis also possible that with the higher inhibitor concen-tration, more complete coverage of the surface by theprotective film occurred.The greater effectiveness of the inhibitor with oxygen

present, i.e., a 20 times reduction in Icorr compared witha 5 times reduction in tests without air for the inhibitorconcentration of 400 ppm, is a strong indication of therole of oxygen in the inhibition mechanism. The higherEpit values obtained in the solutions open to air couldalso be associated with the presence of oxygen. The dataobtained both in the presence and absence of air indicatethat La(4OHcin)3 is acting as an anodic inhibitor. Thedifferences between the Epit and OCP or Ecorr alsoindicate the strength of inhibition. These observationsare consistent with results obtained for mild steel by Blinet al.[7,11] and Behrouzvaziri.[13]

Fig. 14—Scanning electron micrographs of the surface of a tensile specimen tested in 0.01M NaCl with 400 ppm inhibitor at 2 9 10�6 mm/sshowing the fine nodular structure of the cracked background film. The spectrum markers indicate the location of EDS analyses.

-0.8

-0.6

-0.4

-0.2

0

0.2

-6 -5 -4 -3 -2 -1 0 1

Log i mA cm-2

E V

sce b

a

c

Fig. 15—Polarization curves for tests in (a) 0.01M NaCl, (b)0.01MNaCl plus 200 ppm La(4OHcin)3, and (c) 0.01MNaCl plus400 ppm La(4OHcin)3. All solutions open to air.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 43A, JULY 2012—2257

IV. DISCUSSION

The electrochemical data indicate clearly that for SAE4340 steel in NaCl solution, La(4OHcin)3 acts as astrong anodic inhibitor, with the level of inhibitionincreasing with increasing La(4OHcin)3 concentration.The presence of a film with a complex structureconsisting largely of La and Fe oxide and C (as shownby SEM and EDS) is responsible for the electrochemicalresponse and for protecting the steel against corrosionduring the course of the SSRT. The exact nature of theC presence is not known at this time and further studiesare required, although it is likely the C in combinationwith the O originated from the 4-hydroxy cinnamateanion.

It is highly likely that the film is similar to the filmcharacterized by Blin et al.[8,9] in their studies of theinhibition of corrosion on mild steel with La(4OHcin)3.Blin et al.[8] believed the film to be complex chemicallyand to consist of both La hydroxide species andbimetallic complexes. Later studies by Blin et al. usingelectrochemical impedance spectroscopy[7] and infraredspectroscopy[9] suggested that this film consisted of twolayers with the La oxide hydroxide complex forming theouter layer.

Forsyth et al.[12] and Blin et al.[8] proposed thepossible mechanism subsequently for film formation.

(1) The inhibitor molecules bond to the steel surfacethrough the cinnamate moiety, forming a bimetal-lic complex with the iron.

(2) Some corrosion occurs initially with the accompa-nying formation of hydroxyl ions through thecathodic reduction of oxygen.

(3) Some of the La(4OHcin)3 bonds are hydrolyzed bythe local increase in alkalinity to allow complex Laoxide-hydroxide species to form and precipitate.

(4) This process continues across the surface, leadingto a spreading and densification of the film.

In the present study, the formation of a similar film onthe SAE 4340 steel tensile specimens tested in thepresence of La(4OHcin)3 is thought to be responsible forraising the pitting potential, preventing widespreadcorrosion of the steel and the formation of pits. Step 2requires some light corrosion of the steel to occurinitially. It was not possible within the limitations of theSEM used to obtain evidence of this corrosion whenLa(4OHcin)3 was present, although it may well haveexisted. Further surface microscopy is required toidentify the presence of some form of corrosion damageon the specimens tested in the presence of the inhibitor,to help support the proposed mechanism.In the absence of the inhibitor in the SSRT tests,

pitting occurred over the unprotected gage length and inthe notch. It is well known that the development ofpitting in steel is accompanied by the production of lowpH levels within a pit due to the hydrolysis of the Feions and the subsequent production of hydrogen.[36] It isproposed that the hydrogen associated with the pitformation diffused to the regions of high stress triaxi-ality adjacent to pits within the notch and that inter-granular HE cracks initiated (Figure 9) and propagated.The presence of these cracks was responsible for thereduction in the fracture stress compared with the valuesobtained in air. It is beyond the scope of this article todiscuss the mechanisms of HE, but Gangloff classifiedthem in detail under the headings (1) H enhanced bonddecohesion, (2) H adsorption induced localized plastic-ity, and (3) adsorption induced dislocation emission.[17]

Although no obvious signs of plastic deformation wereobserved on the stress vs time plots, some localizeddeformation was observed around the circumference ofthe fracture surfaces containing the intergranular HEcracking. It is possible that localized deformation andthe creation of fresh surface generated during loadingmay have assisted with the uptake of hydrogen into thesteel, as suggested by Scully and Moran.[37]

Table II. Effects of Inhibitor Concentration and Solution Aeration on Various Parameters from Polarization Tests

Condition

Ecorr VSCE Epit VSCE Icorr mA cm�2

Open to Air Deaerated Open to Air Deaerated Open to Air Deaerated

No Inhibitor –0.50 –0.78 — –0.65 2.5 ± 2.4 9 10�3 1.0 ± 0.06 9 10�4

200 ppm Inhibitor –0.33 –0.72 –0.20 –0.24 1.6 ± 0.4 9 10�3 4 9 10�5

400 ppm Inhibitor –0.22 –0.60 –0.10 –0.22 2.9 ± 0.1 9 10�4 2 9 10�5

-1.2

-1

-0.8

-0.6

-0.4

-0.2

0

-8 -7 -6 -5 -4 -3 -2 -1 0

Log i mA cm-2

E V

sce

c

a

b

Fig. 16—Polarization curves for tests in (a) 0.01M NaCl, (b) 0.01MNaCl plus 200 ppm La(4OHcin)3, and (c) 0.01M NaCl plus 400 ppmLa(4OHcin)3. All solutions deaerated.

2258—VOLUME 43A, JULY 2012 METALLURGICAL AND MATERIALS TRANSACTIONS A

The increase in the Epit away from Ecorr and theabsence of widespread pitting and rusting of thespecimens when the inhibitor was present suggests thatLa(4OHcin)3 in preventing widespread pitting alsoprevented accompanying hydrogen production. It isbelieved that the complex film containing La and Feoxides and C observed on the specimen tested in thepresence of the inhibitor was responsible for theprevention of pitting, and subsequent HE of the steel.

V. CONCLUSIONS

1. SAE 4340 steel was shown to be susceptible to HEwhen loaded under slow strain rate conditions in0.01M NaCl solution. The HE is thought to be theresult of corrosion pitting occurring during the testand the accompanying generation of hydrogenwithin the pits.

2. In solution, the inhibitor La(4OHcin)3 caused acomplex film to form on the surface. This film pro-vided anodic inhibition by raising the pitting poten-tial to values higher than the OCP, therebypreventing the formation of widespread pitting.

3. The presence of the inhibitor La(4OHcin)3 at400 ppm in 0.01M NaCl prevented HE from occur-ring, as indicated by no change in the fracture stresscompared to that obtained for tests in air.

REFERENCES1. M. Costa: Crit. Rev. Toxicol., 1997, vol. 27 (5), pp. 431–42.2. B.R.W. Hinton, D.A. Arnott, and N.E. Ryan: Met Forum, 1984,

vol. 7 (4), pp. 211–12.3. B.R.W. Hinton: J. Alloys Compd., 1992, vol. 180 (1–2), pp. 15–25.4. B.R.W. Hinton: Corrosion 89, NACE, New Orleans, LA, 1989,

paper no. 170.5. B.R.W. Hinton, P.N. Trathen, L. Wilson, and N.E. Ryan: 28th

Ann. Conf. of Australasian Corrosion Association, Perth, 1988,Australasian Corrosion Association, Perth, Australia, 1988.

6. A.D.Mercer:Proc. 5th Eur. Symp. on Corrosion Inhibitors, Ferrara,Italy, 1980.

7. F. Blin, P. Koutsoukos, P. Klepetsianis, and M. Forsyth: Elec-trochim Acta, 2007, vol. 52 (21), pp. 6212–20.

8. F. Blin, S.G. Leary, G.B. Deacon, P.C. Junk, and M. Forsyth:Corrosion & Prevention 04 2004, Perth, Australia, Nov. 21–24,2004, Australasian Corrosion Association, Perth, Australia.

9. F. Blin, S.G. Leary, G.B. Deacon, P.C. Junk, and M. Forsyth:Corros. Sci., 2006, vol. 48 (2), pp. 404–19.

10. Corrosion Control & NDT Conf., Melbourne, 2003, F. Blin, S.G.Leary, K. Wilson, G.B. Deacon, P.C. Junk, and M. Forsyth, eds.,Australasian Corrosion Association, Melbourne, Australia.

11. F. Blin, S.G. Leary, K. Wilson, G.B. Deacon, P.C. Junk, and M.Forsyth: J. Appl. Electrochem., 2004, pp. 591–99.

12. M. Forsyth, K. Wilson, T. Behrsing, C. Forsyth, G. Deacon, andA. Phansogoankar: Corrosion, 2002, vol. 58 (11), pp. 953–60.

13. M. Behrouzvaziri, B.R.W. Hinton, and M. Forsyth: Corrosion andPrevention 2008, Wellington, New Zealand, Nov. 2008, Austral-asian Corrosion Association, Australia.

14. M.Behrouzvaziri:MEngScThesis,MonashUniversity,Melbourne,2009.

15. D.R. Gerrard: MEngSc Thesis, Monash University, Melbourne,Australia, 2002.

16. D.R. Gerrard, B.R.W. Hinton, E. Pereloma, and B.W. Cherry:Corrosion and Prevention 1998, Hobart, Nov. 1998, AustralasianCorrosion Association, Blackburn, Australia, 1998, paper no. 38-012.

17. R.P. Gangloff: in Comprehensive Structural Integrity Elsevier Sci-ence, I. Milne, R.O. Ritchie, and B. Karihaloo, eds., ElsevierScience, New York, NY, 2003.

18. R.P. Gangloff, S.A. Shipilov, R.H. Jones, J.M. Olive, and R.B.Rebak: Environment-Induced Cracking of Materials, Elsevier,Amsterdam, 2008, pp. 141–65.

19. C. Barth, E. Steigerwald, and A. Troiano: Corrosion, 1969, vol. 25,pp. 353–58.

20. B.F. Brown: in The Theory of Stress Corrosion in Alloys, J.C.Scully, ed., NATO, Brussels, 1971.

21. W.J. Pollock: Corros. Sci., 1992, vol. 33 (7), pp. 1105–19.22. N. Winzer, A. Atrens, W. Dietzel, G. Song, and K. Kainer:Mater.

Sci. Eng. A, 2008, vol. 472, pp. 97–106.23. S. Ramamurthy and A. Atrens: Corros. Sci., 2010, vol. 52,

pp. 1042–51.24. S. Ramamurthy, W. Lau, and A. Atrens: Corros. Sci., 2011,

vol. 53, pp. 2419–24.25. L. Raymond: in ASM Materials Handbook, J.R. Davis, ed., ASM

INTERNATIONAL, Metals Park, OH, 1987, p. 288.26. Standard Test Method for Mechanical Hydrogen Embrittlement

Evaluation of Plating/Coating Processes and Service Environments,ASTM F519-10. ASTM International, West Conshohocken, PA,2008.

27. W.J. Pollock: 4th Int. Conf. on Hydrogen in Metals, Beijing, 1988.28. Standard Practice for Slow Strain Rate Testing to Evaluate the

Susceptibility of Metallic Materials to Environmentally AssistedCracking ASTM G129-00(2006), ASTM International, WestConshohocken, PA, 2006.

29. Standard Reference Test Method for Making Potentiostatic andPotentiodynamic Anodic Polarization Measurements, ASTM G5-94(2004), ASTM International, West Conshohocken, PA, 2004.

30. G.B. Deacon, M. Forsyth, P.C. Junk, S.G. Leary, and W.W. Lee:Z Anorgan. Allgem. Chemie, 2009, vol. 635 (6–7), pp. 833–39.

31. W.J. Pollock and B.R.W. Hinton: in Galvanic Corrosion, ASTMSTP 978. H.P. Hack, ed., ASTM, Philadelphia, PA, 1988, pp. 35–50.

32. R.J. Sanford: Principles of Fracture Mechanics, Prentice Hall,Englewood Cliffs, NJ, 2003.

33. D. Broek: Elementary Engineering Fracture Mechanics, 4th ed.,Kluwer Academic Publishers, Boston, MA, 1986.

34. G.E. Dieter: Mechanical Metallurgy, 2nd ed., McGraw-HillKogakusha, Tokyo, 1961.

35. M.R. Louthan: in Tensile Testing, 2nd ed., J.R. Davis, ed., ASMINTERNATIONAL, Materials Park, OH, 2004, pp. 115–36.

36. G. Wranglen: An Introduction to Corrosion and Protection ofMetals, Chapman and Hall, London, 1985.

37. J.R. Scully and P.J. Moran: Corrosion, 1988, vol. 44 (3), pp. 176–86.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 43A, JULY 2012—2259