Comparison of the intrinsic properties of EBPVD Al–Ti and Al–Mg coatings

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Materials Chemistry and Physics 132 (2012) 154– 161

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Materials Chemistry and Physics

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omparison of the intrinsic properties of EBPVD Al–Ti and Al–Mg coatings

. Pereza,∗, F. Sanchetteb, A. Billardc, C. Rébéréa, C. Berzioua, S. Touzaina, J. Creusa

Laboratoire d’Etudes des Matériaux en Milieux Agressifs (LEMMA), Université de la Rochelle, Avenue Michel Crépeau, F-17042 La Rochelle Cedex 01, FranceCEA Grenoble DRT/LITEN/DTMN/LTS, 17 rue des Martyrs, F-38054 Grenoble, FranceLaboratoire d’Etudes et de Recherches sur les Matériaux, les procédés et les surfaces (LERMPS)-UTBM, F-90010 Belfort, France

r t i c l e i n f o

rticle history:eceived 18 January 2011eceived in revised form 12 October 2011ccepted 11 November 2011

eywords:l-based coatingslectron-beam evaporative PVDorrosionardness

a b s t r a c t

Aluminium based alloys were deposited on glass slide substrates by electron-beam evaporative PVD(EBPVD) technique in order to assess to the intrinsic corrosion behaviour in saline solution. Alloying ele-ments (Ti or Mg) were added in order to improve both mechanical and corrosion behaviours of aluminiumfilms. The goal of this study is to develop new alternative sacrificial coatings that could protect steel flatproducts against corrosion with enhanced mechanical properties in order to limit the seizing sensitivity.Because of the difference in vapour pressures, aluminium and alloying elements were co-evaporatedfrom two separate sources. Mixing of the vapour flux was the most appropriate experimental method toensure a homogeneous distribution of the alloying element into the aluminium based coating. Differentcontents of transition metals were examined, and it was noticed that the microstructure evolves from asupersaturated solid solution to an amorphous phase due to the lattice distortions with the incorporation

of the alloying elements. Electrochemical behaviour of the EBPVD coatings in saline solution was inves-tigated. It is shown that aluminium can be mechanically reinforced while preserving sacrificial corrosionprotection properties. The evolution of the properties was found to be strongly linked to the nature ofthe incorporated element. A comparison with a previous study on the electrochemical and mechanicalbehaviour of EBPVD Al–Cr and Al–Gd coatings permitted a comprehensive analysis of the influence ofthe incorporation of transition metals in aluminium.

© 2011 Elsevier B.V. All rights reserved.

. Introduction

Steels can be protected against corrosion by protective coat-ngs like zinc, chromium or nickel based coatings [1,2]. However,he environmental problems due to the elaboration of these coat-ngs through electrolytic process lead to the research of alternativeolutions. Aluminium coatings elaborated by physical vapour depo-ition (PVD) are a good choice but they tend to passivate, they areensitive to localized corrosion in saline environments [3–5] andave poor mechanical properties. Monolayers have been tested

n order to find alloying elements that will change the corrosionehaviour and/or will enhance mechanical properties [6–15]. The

ncorporation of transition metals like Cr, Mn and Mo improveshe mechanical properties of coatings, but induces an ennoble-

ent of the corrosion potential and thus reduced the sacrificialharacter of the alloy film [16–20]. By incorporation of metals less

oble than aluminium like Gd, Nd, and Sm the sacrificial characteran be preserved, but the mechanical properties are not signifi-antly improved [10,13,14]. Most of these studies are focused on the

∗ Corresponding author. Tel.: +33 5 46 45 72 94; fax: +33 5 46 45 72 72.E-mail address: perez.andrea@laposte.net (A. Perez).

254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.matchemphys.2011.11.013

intrinsic corrosion behaviour of Al based films, deposited by con-ventional magnetron sputtering techniques onto an inert substrate.Electrochemically inert substrates (glass slides, silicon wafers orvitreous carbon) are widely chosen in the literature [9,21–25] inorder to assess to the intrinsic behaviour of the deposited filmswithout interaction with the substrate. Such a step is favourableand largely widespread to predict the risks of galvanic couplingthanks to the superposition of the polarization curves of the coat-ings and the substrate, simulating different surface ratio [26].

Furthermore, many studies examine the behaviour of a uniquealuminium alloy deposited by conventional sputtering techniqueseither evaluating the influence of the alloying element content,or investigating the anodisation behaviour in electrolytes of con-trolled composition [2,12,13,23,24]. In fact [2,6–11], Al–Ti andAl–Mg alloys, deposited by magnetron sputtering, were foundrespectively to reinforce the mechanical properties and enhancethe corrosion resistance. Only few studies are focused on the com-parison of the mechanical and/or electrochemical properties ofdifferent Al alloys deposited by electron-beam PVD (that provides

significant deposition rates). The EBPVD process seems to be moreappropriate for industrial applications, offering independence andfreedom in control of microstructure and composition of the coat-ing via manipulation of controlling process parameters such as

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the amorphous phase may occur in the range between 24 and32 at.% Ti. However, for magnetron sputtered alloys, the amorphousphase appears between 25 and 60 at.% [36]. This difference maybe linked to the less dense alloys formed by EBPVD than the ones

Table 1Composition of the different samples, with the alloying element at.% composition.

Sample

A. Perez et al. / Materials Chemi

ressure level, substrate temperature, power (voltage, current) andvaporant composition. This technique allows the production oflms from a wide range of materials such as, metallic, dielectricsnd semiconductors with relatively low deposition time [27].

EBPVD film growth mechanisms remain the same as those ofoatings deposited by other CVD or PVD techniques, including sput-ering. According to thin film nucleation and growth mechanismsescribed in the literature [28,29], the final structure of the film

s a result of the creation of multiple nucleation events followedy subsequent growth of the structural units (grains in crystallineaterials) and their periodic reorganization during coalescence.

he growth and coalescence processes are directly dependent onubstrate deposition temperature and total energy of the system. Sot can be considered that individual grains grow in the form of sin-le crystals. Although structure zone models can explain the variousicrostructures which occur in PVD coatings as a function of both

hermal and ion induced atom movements, no account is taken onhe influence of the substrate on the microstructure of the coatings.he substrate crystal structure can affect the growth of the deposithrough changes in the distribution, orientation, and size of therystallites nucleated on the growing surface and the nature of theubstrate can also affect the level of internal stresses in a film due toifference in thermal expansion coefficients [30]. Recent studies onhe influence of the metallurgical state evolutions (crystallographicrientations, grain size, deformation, etc.) on the reactivity of fccetals [31,32] revealed that the degradation mechanisms were notodified, only the kinetics were affected. According to these stud-

es, substrate changes during the PVD process could effectively alterhe microstructure of the coatings, but the potential domain whereeposit degradation occurs is considered to remain unmodified.his fact is often taken into account when the intrinsic behaviourf PVD coatings is investigated [9,24,25].

This study presents the characterization of morphology,icrostructure and properties for different compositions of Al–Ti

nd Al–Mg coatings deposited by an electron-beam evaporativeVD technique. The alloy coatings are studied on glass slides, inrder to evaluate their intrinsic properties and then simulate thealvanic coupling between the coating and the steel.

. Experimental procedure

Al–Ti and Al–Mg coatings were deposited as monolayers on glass slides in aLASSYS EBPVD system equipped with 2 electron beam evaporators. EB-PVD is aimple process in which a focused high energy electron beam is directed to melt thevaporant material(s) in a vacuum chamber. The evaporating material condenses onhe surface of the substrates or components resulting in the formation of deposit, i.e.,oating. Experimental procedure for coatings preparation is the same as for Al–Crnd Al–Gd described elsewhere [10]. Substrates were ultrasonically cleaned in ace-one, rinsed in alcohol and etched in the vacuum chamber before the deposition stepy means of an argon ion gun (500 eV). Deposition rate was in the range 1–4 nm s−1.ll coatings, 5 �m thick, were deposited at low temperature (below 150 ◦C) and lowressure, about 10−4 Pa. The substrate–crucible distance was 390 mm.

Al–Ti or Al–Mg monolayers cannot be evaporated simultaneously from the samerucible because, as for Al–Cr and Al–Gd [10], vapour pressures for both liquid com-onents are not within a factor of 100 of each other for a given temperature. Thisriterion is met, for example, for aluminium–iron alloys. In the case of aluminiumnd magnesium, the difference between vapour pressures is very high but a goodontrol of composition can be achieved. Nevertheless, the control of the coating com-osition must be carried out carefully via separate control of evaporation rates. Thisechnique also leads to some unusual levels of supersaturation of both interstitialnd low-miscibility substitutional alloying elements [33]. Hardness was measuredsing a Nano-Hardness Tester (CSEM NHT) with a Vickers indenter tip. 40 inden-ations were made in each sample and the results presented are extracted fromtatistics on indentation results. Hardness was determined using the Oliver andharr analysis [34].

The surface morphology, before and after polarization testing, was examined

y optical microscopy using a LEICA DM6000M and the average composition of theluminium based alloys was determined, at different scales along the glass slidey microanalyses with an energy dispersive X-ray (EDX) spectrometer coupled to

scanning electron microscope (SEM) Micro Quanta 200 SEM/FEG FEI Phillips. Theicrostructure of the coatings was examined by X-ray diffraction (XRD) using a

d Physics 132 (2012) 154– 161 155

Bruker Advance D8 diffractometer with Cu K� radiation and a step size of 0.02◦ inthe range of 20–120◦ of two-theta.

The immersion tests were performed in an aerated and stirred NaCl 5 wt.% solu-tion with pH adjusted to 7.0 by addition of diluted NaOH solution and at temperatureof 25.0 ± 0.1 ◦C. The electrochemical experiments were carried out in a conventionalthree electrode glass cell using samples as working electrode, connected to a Bio-Logic VSP potentiostat and piloted by EC-lab software. The potential is referred to asaturated calomel electrode (SCE) and the counter electrode was a large platinumgrid. Polarization curves were obtained after an initial potential stabilization of 1 h inthe solution, into a range of ±150 mV around the open circuit potential (OCP), and ascan rate of 0.2 mV s−1. A current density threshold of 100 �A cm−2 was fixed in orderto evaluate the sensitivity to localized corrosion in saline solution. The corrosionpotential Ecorr and current density icorr were estimated by using Tafel extrapola-tion [35]. Immersion tests during 48 h were also performed with OCP monitoringand periodic polarization resistance (Rp) measurements. These last experimentswere performed using linear polarization tests consisting of a potential sweep of±20 mV per OCP at a scan rate of 0.2 mV s−1. All electrochemical measurementswere repeated twice or more to verify the reproducibility.

3. Results and discussion

The average composition of the samples obtained by EDX anal-ysis is given in Table 1, with only the percent of the incorporatedalloying element. It can be underlined that the variation of the alloy-ing element content along the axis of the glass slide was found tobe very low. So, it is considered that the distribution of the alloyingelement is homogeneous for all the samples.

Fig. 1 presents the XRD diagrams for the Al–Ti coatings (Fig. 1a)and for the Al–Mg coatings (Fig. 1b). For titanium or magnesiumcontents up to respectively <24 at.% Ti and <4 at.% Mg, the singlephase coatings have a fcc crystalline structure corresponding toa supersaturated �-aluminium solid solution of Ti or Mg in alu-minium. The lattice parameter estimated from the angular positionof the (1 1 1) Al peak does not change significantly with the incor-poration of titanium whereas a slight increase is detected withthe incorporation of magnesium. It is due to similar atomic radiibetween aluminium and titanium and a larger size of magnesium(0.146 nm for titanium, 0.160 nm for magnesium and 0.143 nmfor aluminium). This classical result is consistent with previousstudies whatever the PVD technology used [7,10]. As previously,the solubility limits are different because the effect of differencesbetween the size of their atomic radius and that of the aluminiumatom. Addition of titanium or magnesium in �-aluminium phasefavours the (1 1 1)�-Al texture. It is noticeable that single-phased �-titanium alloys exhibit a strong (0 0 2)�-Ti texture (32 at.% in Fig. 1a).This particular effect is expected to proceed from the close-packedstructure of aluminium and titanium [36]. In all cases, diffractionpeak broadening is observed, which could be either associated to adecrease of the mean �-aluminium grain size [37] or to an increaseof the internal stresses. However, the shift of the peaks is not suffi-cient to explain the peak broadening by an increase of the internalstresses.

Unlike for sputtered Al–Cr, Al–Ti, Al–Fe [24] or for e-beam evap-orated Al–Gd alloys [10], no amorphous phase was detected byX-ray diffraction in the Al–Ti system (Fig. 1a). We can assume that

1 2 3 4 5 6

Mg (at.%) 3 4 9 16 18 21Ti (at.%) 2 9 17 19 24 32

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156 A. Perez et al. / Materials Chemistry and Physics 132 (2012) 154– 161

Fig. 1. XRD diagrams of the Al–Ti alloys (a) and the Al–Mg alloys (b).

Fig. 3. Hardness of as-deposited Al–Ti and Al–Mg coatings.

Fig. 2. Scanning electron micrographs of fracture cro

formed by magnetron sputtering. For the 32 at.% alloy, the coatingseems to be a metastable solid solution which derives from hcp �-Ti whereas, for high magnesium contents in Al (i.e. from 9 to 21 at.%Mg), the coatings are a mixture of amorphous and crystallized �-Al phases (Fig. 1b).Regarding the morphology of the coatings, theyare generally less compact than the same binary alloys obtained bymagnetron sputtering [7,36]. Cross-sections of different samplespresented in Fig. 2 clearly show columnar morphologies (whateverthe alloying element contents) in both Al–Ti and Al–Mg systems.Nevertheless, a slight densification for higher contents could beattributed to the crystalline growth associated with the decrease ofthe mean grain size [37,38]. Unlike amorphous single phase alloysdeposited by magnetron sputtering, no fully compact morphologywas observed on fracture cross-section of electron beam evapo-

rated coatings. Indeed, in this case, the well known effect of atomicpeening and grit blasting during films growth cannot be invoked,due to the non-conductive glass substrates.

ss-sections of Al–Ti (a) and Al–Mg (b) deposits.

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Fig. 3 shows the influence of the alloying element content onano-hardness. Hardening is observed for both titanium and mag-esium addition in aluminium. This phenomenon is likely to beue to a combination of a solid solution effects and grain sizeefinement. Nevertheless titanium addition has a lower effect whenoatings are deposited by e-beam evaporative technique comparedith as-deposited magnetron sputtering films: 7 GPa for 25 at.% Ti

7]. This result is the same for the Al–Mg alloys: for 7 at.% Mg, theardness is 4.5 GPa [15]. It is probably due to a less compact mor-hology associated with the lower energy of impinging atoms onhe growing film with the EBPVD process mainly because of a veryow ionization level, and of a low working temperature. It is alsohown in Fig. 3 and with the previous results on Al–Cr and Al–Gdystems [10] that the increase in hardness of the �-Al solid solutions more pronounced in the case of the addition of element whosetomic radius is widely different from Al one. Thus, �-Al(Gd) islightly harder than �-Al(Mg) which is harder than �-Al(Ti) and-Al(Cr) solid solutions.

For high alloying element contents, hardness values graduallyncrease with titanium incorporation, probably due to modifi-ation of the microstructure and the presence of a metastablec-Ti phase. For magnesium incorporation, the presence of aimodal structure at high Mg content containing an amor-hous phase seems to limit the values of hardness, as waslready observed for other PVD coatings with amorphous phaseonstituents [6,9].

Fig. 4 shows the polarization curves of the different samplesfter 1 h of immersion in saline solution. All the Al–Ti alloys (Fig. 4and b) present a large passive domain and localized effects asso-iated with pitting corrosion. Polarization curves can be comparedo the corrosion region of the steel substrates (symbolised by theatched oblong in Fig. 4). All the Al–Ti alloys are sacrificial againstteel (Fig. 4a and b), but except for the 2 at.% content, their pittingotential is in the corrosion region of the steels. Thus, these alloysill not keep a sacrificial behaviour during a long time due to the

alvanic coupling between steel and the coating. The 2 at.% alloy has pitting potential lower than the corrosion region of steel, howeverery close. So, it can be assumed that the protection provided willot be efficient due to a low cathodic polarization when coupled toteel.

The Al–Mg alloys also present a passive domain, followed byocalized corrosion. Comparison of the polarization curves of theamples and the corrosion region of steels shows that, for all theompositions in magnesium, the alloy is sacrificial compared toteels. But very negative values of corrosion potential may leado the hydrogen evolution reaction (HER) and to embrittlement ofhe steel (the threshold potential for HER in aluminium is about1.2 V per SCE) [39]. So it is important to avoid such compositionomains that would be detrimental for the mechanical propertiesf the coated steel. A composition that would be appropriate forteel protection is 3 or 4 at.% Mg.

The electrochemical characteristics of the different Al–Ti sam-les are shown in Fig. 5. Each point is representative of several tests.he general trend for corrosion potential of Al–Ti alloys seems toe an enoblement of the potential with titanium incorporation upo 19 at.%, and then a stabilization. The pitting potential increasesinearly with the titanium content (Fig. 5b): the incorporation ofitanium seems to stabilize the passivation of the alloy, whichecomes more resistant to localized corrosion. This behaviour isimilar to the one of Al–Ti alloys deposited by magnetron sput-ering [6]. The ennoblement of the corrosion potential limits theathodic protection provided by these coatings. The shift of the

itting potential towards more positive values with Ti content isarmful for the sacrificial character during the ageing of a coatedtructure. The EBPVD Al–Ti alloys do not present electrochemi-al properties that could be applied for the cathodic protection

Fig. 4. Polarization curves of the Al–Ti alloys (a and b) and of the Al–Mg alloys (c)after 1 h of immersion in saline solution.

of steels, and then may not be chosen as a sacrificial coating ina multilayer configuration.

The ennoblement of the pitting potential can first be explainedwith the passive film composition. Some authors showed that tita-nium oxide is incorporated in the passive film as well as aluminiumoxide [40]. The two oxides are present in the alloy proportions inthe passive film. This observation indicates that titanium is a pas-sivity promoter [41]. As a consequence, it will ennoble the pittingpotential. Moreover, the passive film composition of titanium alloywill affect the pH at the point of zero charge (PZC) [42], pHPZC. Thispoint defines the pH where the surface charge is equal to zero. For

pH lower than pHPZC, the surface has a global positive charge. Theadsorption of chloride ions is easier and facilitates breakdown of thepassive film. The pHPZC of titanium oxide is included between 4.7

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ig. 5. Evolutions of the corrosion potential and the corrosion current density verurves obtained after 1 h of immersion in saline solution.

nd 6.7 whereas the one of aluminium oxide is included between.6 and 7.7 [43]. The pHPZC of the passive film depends on the Tiontent and we assume that it decreases with the incorporationf titanium to come closer to the one of the titanium oxide. Thus,t can be proposed that with a lower pHPZC, less chloride ions aredsorbed on the surface of the Al–Ti passive film and the pittingesistance is therefore increased.

The corrosion behaviour of the Al–Ti alloys can be compared tohe one of the Al–Cr alloys [10]: the corrosion potential increasesith the content of alloying element and the pitting resistance too.n the other hand, the current corrosion density stays in the same

ange of values when Ti content increases whereas it increases withr content. So Ti stabilizes the passivation of the coating withouthanging drastically its reactivity while Cr stabilizes the passivationnd increases the reactivity.

Concerning the Al–Mg alloys, their corrosion potential is alwaysower than −1.1 V per SCE. The trend for the corrosion potentialvolution is a decrease when magnesium content increases. Fromround 15 at.% Mg, the HER becomes predominant. It can affecthe corrosion mechanisms and induce a variability of the corrosionotential.

This evolution is linked to the one of the corrosion currentensity, which follows the same evolution at around the sameagnesium contents. This could be explained by the corrosion

roducts, which cover the surface of the sample and then act as

e Ti (a and b) and versus the Mg (c and d) content deduced from the polarization

a physical barrier. However, the pitting potential does not evolvesignificantly. This behaviour is slightly different from the one ofAl–Mg alloys deposited by magnetron sputtering [2,42], whose cor-rosion potential decreases when magnesium content increases, andwhose corrosion current density increases with magnesium con-tent [44]. This difference may be due to a morphology less compactwith EBPVD alloys than in PVD magnetron alloys.

The corrosion behaviour of these alloys can be compared to theone of Al–Gd alloys [10]: the corrosion potential decreases whilethe alloying element content increases, whereas the corrosion cur-rent density increases. The reactivity increases with the alloyingelement content. The values of the pitting potentials differ: theone of the Al–Mg does not change significantly whereas the oneof the Al–Gd decreases when gadolinium content increases. Thus,the passivation of Al–Gd alloys is less stable than Al–Mg alloys.

Immersion tests have been conducted for 48 h on three compo-sitions Al–Mg (3 at.%, 9 at.%, 18 at.%). The OCP curves are presentedin Fig. 6. The Al–Mg 18 at.% presents a very significant reactivity, andthe presence of a large pit induces infiltration of solution beneaththe coating. So the coating was detached from the glass slide andthe experiment stopped. Some transient stages can be observed at

the beginning of the immersion. These transient stages are linkedto polarization resistance measure (at 30 m, 1 h and 2 h). Actu-ally, this measures implies a slight polarization of the coating andthe magnesium being very reactive, the potential can be changed

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Table 2Average ratio Al/Mg calculated with the EDS measurements.

Surface beforeimmersion

Surface afterimmersion,outside a pit

Surface afterimmersion,inside a pit

Fig. 6. Open circuit potentials for AlMg 3 at.%, AlMg 9 at.% and AlMg 18.%.

fter the measure. The polarization resistances are not representedere because of oscillations. Anyway, when magnesium is incorpo-ated, it seems that reactivity is increased. As observed for EBPVDl–Gd alloys [10], this phenomenon is related to the large cathodichift of the corrosion potential that allows the hydrogen evolutioneaction as presented in Fig. 6. Only alloys with content less than0 at.% of Mg present interesting reactivity that could be applied aslternative candidates for cathodic protection of steels, which wasifferent from results obtained from magnetron sputtering. Actu-lly, Baldwin [2] reports that the most efficient sacrificial coatings the one with 20 wt.% Mg (22 at.% Mg). The high reactivity coulde attributed to a selective dissolution of magnesium comparedo aluminium as reported in Fe–Mg alloys [45], so the composi-ion on the surface may evolve. But as the alloys tend to passivate,he selective dissolution may be prevented. If a modification of theomposition occurs, it would be localized inside the pits favouredy the acidification of the solution during the pit growth.

EDX measurements have been performed on the surface of theample after immersion and the Al/Mg ratio was calculated insidehe pit, outside the pit and on the sample and compared to the

atio value on the surface before immersion in order to see whetherr not there is a preferential magnesium dissolution. All the pitsave the same morphology, presented in Fig. 7. The average ratio

Fig. 7. SEM picture of a pit on the surface of the AlMg 9 at.%.

Ratio Al/Mg (meanon 3 values)

9 9 12

Al/Mg calculated by EDX measurements is presented in Table 2. Theratio increases inside the pits, thus indicating that either the alu-minium content increases; either that of the magnesium contentdecreases. The second information given by the EDX measurementsis that the ratio Al/Mg is the same before immersion and afterimmersion outside the pit. The passive film formed during the firststage of immersion prevents any selective dissolution of magne-sium compared to aluminium. There is a preferential dissolution ofthe magnesium localized only in the pits essentially due to acidi-fication of the solution that accelerates the dissolution of coating,and the compositions of the surface outside the pits do not changesignificantly.

Three Al–Ti alloys compositions were also tested for 48 h insaline immersion: Al–Ti 2 at.%, Al–Ti 9 at.%, Al–Ti 17 at.%. A decreaseof OCP (Fig. 8) is observed at the beginning of the immersion,followed by a stabilization of the potential. For Al–Ti 2 at.%, thepolarization resistance (Rp) increases quickly at the beginning ofthe immersion and then stabilizes (Fig. 9). For Al–Ti 9 at.% and17 at.%, the polarization resistance (Rp) decreases then stabilizes.A decrease of both potential and Rp can be linked to pit forma-tion on the surface of the samples. After the formation of pits, thepotential remains constant at the pitting potential of the samples.For the Al–Ti 9 at.%, the potential decreases again at about 30 h.This is linked to the numerous pits formed on the surface of thesample, which leads to more uniform corrosion. The potential is inconclusion more close to the corrosion potential observed duringthe polarization. The low values of Rp indicate a high reactivity ofthe alloys 9 and 17 at.%. However, for Al–Ti 2 at.%, the high valuesof Rp (500 k� cm2) suggests a very low reactivity. A lot of pits areformed on the surface of the sample so the low reactivity suggeststhat the pits are metastable and repassivate very quickly.

Some statistics have been calculated on the pits that are presenton the sample surface after immersion, in order to evaluate if a

higher reactivity at the beginning of the immersion could lead tolarger pits or more pits formed on the surface. Fig. 10 shows thedensity of pits classified as function of their diameter. All Al–Ti

Fig. 8. Open circuit potentials for AlTi 2 at.%, AlTi 9 at.% and AlTi 17 at.%.

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160 A. Perez et al. / Materials Chemistry an

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Fig. 9. Polarization resistance for AlTi 2 at.%, AlTi 9 at.% and AlTi 17 at.%.

lloys present very numerous pits with diameters of between 10nd 20 �m and the mean diameter of the pits are very similaretween the samples (respectively 15, 25 and 21 �m for the 2 at.%,

at.% and 17 at.%).The 2 at.% sample present a lot of pits on its surface

162 pits cm−2). This alloy presents the lowest reactivity and has themallest pits on its surface, which is consistent with our hypothesisf metastable pits that repassivate. Actually, the OCP measured forhis sample is close to its corrosion potential and it confirms therevious assumption.

When titanium content increases (9 at.%), the pit density is4 pits cm−2. This is linked to the fact that titanium is a passivityromoter and its incorporation into aluminium leads to an eno-lement of the pitting corrosion resistance. So the pit initiation iselayed. Moreover, the pits observed are probably metastable pitsue to their small size. The OCP measured for this sample is, as forhe 2 at.%, its corrosion potential.

Finally, the 17 at.% presents the highest pit density. But the OCPeasured is close to the pitting potential of the sample, and we can

hus assume that the pits formed are stable. So it is not possibleo compare the behaviour of the samples with 2 and 9 at.% withhe one with 17 at.%. Grain size could also play a role for pit den-ity, as was revealed that the passive film defects are linked to therain boundaries [46]. So if grain size decreases, the number of grain

oundaries increases and consequently the number of passive filmefects.

ig. 10. Histogram of the pits on the surface of the Al–Ti samples after the immer-ion.

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d Physics 132 (2012) 154– 161

Preliminary results (not shown here) revealed that the alu-minium presents only few pits but pretty large, whereas the Al–Tialloys present very numerous pits but very small. Statistics onAl–Mg alloys surface has not been calculated, because the pits arevery similar to pits formed on aluminium.

4. Conclusions

Incorporation of transition metals is an adapted method tomodify properties of Al EBPVD coatings. The properties of EBPVDcoatings with different Al–Ti and Al–Mg compositions were stud-ied. The coatings are generally less compact than the same coatingsdeposited with magnetron sputtering, and their morphology ofgrowth is columnar. Besides, both Ti and Mg have a hardeningeffect, but lower than the one of magnetron sputtered coatings.

All the studied samples present a localized corrosion associatedto a passive domain and/or a pitting corrosion.

The Al–Mg alloys, with high Mg content (above 10 at.%) havea too high reactivity and a corrosion potential too negative to beconsidered in a multilayer configuration, because of the HER. Butlow Mg content (below 10 at.%) can be considered in a multilayerconfiguration for sacrificial properties.

Concerning the Al–Ti alloys, incorporation of titanium seems tostabilize the passive domain, which increases the pitting resistanceof the material. The pits are smaller but numerous on the surfaceof the alloys after immersion tests (between 74 and 188 pits cm−2).So Al–Ti alloys effectively present reinforced corrosion behaviourin saline solution, but these electrochemical characteristics are notconsistent with cathodic protection of steel. However, the hard-ening effect of the incorporation of titanium makes them goodcandidates for the layer with good mechanical properties. Thesealloys can be included in a multilayer configuration with a secondlayer presenting interesting electrochemical properties, as Al–Mgwith low Mg content.

However, complementary tests on these coatings depositedonto steel substrates are necessary to validate our conclusions. Thechange of substrate nature could induce modification of morphol-ogy that affects the electrochemical properties essentially in termof degradation rate evolution.

Acknowledgements

This research was supported by a grant from the FrenchResearch National Agency through the national program on Materi-als and Processes (06MAPR0020). The Common Centre of Analyses(C.C.A.) and Egle Conforto are particularly thanked for the technicalsupport and advices in this program.

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