Surface & Coatings Technology 206 (2012) 3903–3910

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Chromium and tantalum oxide nanocoatings prepared by filtered cathodic arcdeposition for corrosion protection of carbon steel

Belén Díaz a,b, Jolanta Światowska a,b, Vincent Maurice a,b,⁎, Marcin Pisarek a,b, Antoine Seyeux a,b,Sandrine Zanna a,b, Sanna Tervakangas c, Jukka Kolehmainen c, Philippe Marcus a,b,⁎a Chimie ParisTech, Laboratoire de Physico-Chimie des Surfaces (LPCS), 11 rue Pierre et Marie Curie, F-75005 Paris, Franceb CNRS UMR 7045, 11 rue Pierre et Marie Curie, F-75005 Paris, Francec DIARC-Technology Inc., Kattilalaaksontie 1, 02330 Espoo, Finland

⁎ Corresponding authors at: Chimie ParisTech, LaboSurfaces (LPCS), 11 rue Pierre, et Marie Curie, F-75005 P

E-mail addresses: vincent-maurice@chimie-paristechphilippe-marcus@chimie-paristech.fr (P. Marcus).

0257-8972/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.surfcoat.2012.03.048

a b s t r a c t

a r t i c l e i n f o

Article history:Received 24 November 2011Accepted in revised form 15 March 2012Available online 6 April 2012

Keywords:Corrosion protectionThin filmsChromium oxideTantalum oxideFCADToF-SIMS

Combined analysis by Time-of-Flight Secondary Ion Mass Spectrometry (ToF-SIMS), X-ray PhotoelectronSpectroscopy (XPS), polarization curves and Electrochemical Impedance Spectroscopy (EIS) of the relationbetween chemical architecture of thin (10 and 50 nm) chromium and tantalum oxide coatings grown byfiltered cathodic arc deposition (FCAD) on carbon steel and their corrosion protection properties is reported.Pre-etching in the deposition process allows reducing the substrate native oxide layer to traces of iron oxide.A carbidic interlayer is then formed by reaction between the first deposited metallic particles and the residualcarbon surface contamination of the alloy. The bulk coatings mostly consist of Cr2O3 or Ta2O5 with no in-depth variation of the stoichiometry. Surface and bulk of the coatings are contaminated by hydroxyl andorganic groups. The 50 nm coating has a relatively large porosity assigned to a columnar growth preventinggood sealing at grain boundaries. The duplex structure (Ta/Ta–C) of the carbidic interlayer promotes a lessdefective growth of tantalum oxide than the single Cr–C interlayer for chromium oxide, thereby improvingthe sealing properties. The dielectric constants suggest poor insulating properties in line with a defective andporous nanostructure of the coatings. No dissolution was observed for both oxide nanocoatings in neutral0.2 M NaCl. Penetration of the electrolyte and access to the interface with the carbon steel surface cause thedissolution of the Cr–C interlayer, but not that of the Ta/Ta–C interlayer, and a more rapid initiation oflocalized corrosion.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

Conventional Physical Vapor Deposition (PVD) techniques such asmagnetron sputtering, ion plating, anodic or cathodic arc depositionare employed to produce corrosion resistance coatings, typically of TiN,TiAlN and CrN [1–3]. Among them cathodic arc deposited coatingsexhibit improved adhesion to the substrate thanks to ion bombardmentprior to deposition that allows reducing chemical heterogeneities androughness to a near atomic level [4,5]. The cathodic arc technique is thusable to produce high quality hard coatings which offer improvedmechanical properties [6] and suitable corrosion protection [7]. Themain disadvantage of this process is the generation of macroparticlesoriginatingmainly from droplet formation during arc evaporation [4,7],which limit structural homogeneity with consequences for the range ofapplications. Several attempts were made for eliminating macroparti-cles.Wang et al. [8] achieved an improved surfacemorphology by using

ratoire de Physico-Chimie desaris, France..fr (V. Maurice),

rights reserved.

filtered cathodes, shifting the distribution of particles to lower sizes.This filtered systemhas been shown to be suitable for the preparation ofhigh quality optical and dielectric materials and conventional wearresistant TiN films [9]. An improved corrosion resistance was alsomeasured with filtered TiAlN coatings deposited on carbon steel, incomparison to unfiltered analogs [10,11].

Several groups have focused their studies on the electrochemicalcharacterization of filtered cathodic arc deposition (FCAD) coatings,commonly in the μm range in thickness. New materials such as TiSiN,Ta–C:N and a-C:Crwere prepared with improved deposition conditionson substrates susceptible to corrosion. The FCAD films were proven toprovide a suitable corrosion resistance by decreasing the corrosion rates[10–14]. The pore density was shown to be the key parameter to bereduced in order to improve the corrosion resistance performance.

Metal oxides such as TiO2, ZrO2, Al2O3, Nb2O5 and SnO2 alsorepresent another group of important FCAD materials due to theirmechanical and chemical stability. The flexibility in varying the opticalproperties gives these metal oxide films a wide range of applications inoptics, microelectronics, optoelectronics as well as in sensors andprotective coatings [7]. Here we report on the characterization of thin(10 and 50 nm)metal oxide layers deposited by the FCAD technique on

3904 B. Díaz et al. / Surface & Coatings Technology 206 (2012) 3903–3910

carbon steel substrates in order to assess their corrosion protectionproperties. Single films of tantalum oxide and chromium oxide weredeposited and studiedwith surface analytical techniques, Time of FlightSecondary Ion Mass Spectrometry (ToF-SIMS) and X-ray PhotoelectronSpectroscopy (XPS), in order to characterize their surface, bulk andinterfacial chemical composition. Polarization curves were employed tocalculate the uncoated surface fraction or so-called coating porosity andElectrochemical Impedance Spectroscopy (EIS) was employed in orderto test the corrosion resistance during immersion in a corrosive sodiumchloride aqueous solution.

2. Experimental

The nominal chemical composition of the hardened (805 HVhardness) and tempered (180 °C temperature) low alloy carbon steel(AISI 52100, DIN 100Cr6) substrate employed in this study is given inTable 1. The surfacewas lappedwith awater based diamond suspension(~6 μm). Prior to coating the substrates were carefully wiped withacetone, ultrasonicated for 5 min in isopropanol, and blow-dried withcompressed air.

The coating process was carried out in DIARC FCAD coatingequipment. Samples were pre-etched with 350 eV Ar ions at a currentdensity 0.5 mA cm−2 for 30 min before coating deposition. The metaloxide coatings Ta2O5 or Cr2O3 were produced by depositing Ta or Crplasma, respectively, from solid targets having min 99.8% purity inpartial pressure of oxygen. The deposition rate of the metal oxidecoatings was approximately 0.5 μm h−1. During deposition thesamples are heated up to a maximum temperature of 50 °C. Thecoating thickness was controlled from a silicon sample which wasfixed in the deposition chamber at a similar position as the steelsamples. The thickness was measured with a Dektak 3ST profilometerfrom the interface between a coated and an uncoated region of thesilicon sample after the removal of ink mask. This procedure wassimilar to that adopted in previous deposition studies on the samecarbon steel substrate and proved reliable [15–18].

A ToF-SIMS 5 spectrometer (IonTof) operating at a pressure of10−9 mbar was used for elemental depth profiling the samples. Thespectrometer was run in the HC-BUNCHEDmode with optimummassresolution but poor lateral space resolution. A pulsed 25 keV Bi+

primary ion source was employed for the analysis, delivering ~0.8 pAof current over a 100 μm×100 μm area. It was interlaced with a 2 keVsputtering Cs+ beam giving an ~80 nA target current over a400 μm×400 μm area. Data acquisition and post-processing analyseswere performed using the Ion-Spec software. The profiles wererecorded with negative secondary ions, more sensitive to fragmentsoriginating from oxide matrices. Before the analysis, all samples werecleaned with ethanol for 10 min in an ultrasonic bath and dried withcompressed air. ToF-SIMS depth profiling was performed on pristinesamples and on samples submitted to the 6 hour corrosion testdescribed below.

X-ray Photoelectron Spectroscopy (XPS) chemical analysis wascarried out using a VG ESCALAB 250 spectrometer operating at aresidual pressure of 10−9 mbar. An Al Kα monochromatizedradiation (hν=1486.6 eV) was employed as an X-ray source. Surveyand high resolution spectra were recorded with pass energies of 100and 20 eV, respectively. The photoelectrons were collected at a 90°take-off angle with respect to the substrate surface. The dataprocessing was performed with the Avantage software using a Shirleybackground. The measured core level intensities were converted into

Table 1Chemical composition (w/w %) of the carbon steel substrate employed in this study.

C Si Mn P S Cr Ni Cu Fe

0.9–1.05

0.15–0.35

0.25–0.45

Max.0.03

Max.0.025

1.35–1.65

Max.0.3

Max.0.3

Balance

atomic concentrations using Scofield values for the photoionizationcross sections σ and calibrated values for the transmission factor T ofthe spectrometer. Pristine samples were analyzed after cleaning asdescribed above.

The electrochemical measurements were performed in a conven-tional three-electrode cell using a potentiostat/galvanostat AutolabPGSTAT30. A Saturated Calomel Electrode (SCE) was employed as thereference electrode and a platinum wire as the auxiliary electrode.The working electrode area was 0.44 cm2. All the experiments wereperformed at room temperature in a 0.2 MNaCl aqueous solution (pH 7)prepared with ultra pure water (resistivity>18 MΩ cm) and reagentgrade chemicals (NaCl Analar Normapur analytical reagent, VWR®BDHProlabo®). The electrolyte was bubbled with Ar for 30 min before testand during experiments.

Potentiodynamic polarization curves were obtained after an initialperiod of 1 h at Open Circuit Potential (OCP). The samples werepolarized from −0.9 V in the anodic direction stopping the measure-ment when the current density was above 1 μA cm−2. The scan ratewas 1 mV s−1 for consistency of examination with other oxidenanocoatings deposited by atomic layer deposition (ALD) on thesame carbon steel [15–18]. The coating efficiency was evaluated fromthe polarization resistance (Rp) values obtained by Tafel analysis aspreviously detailed [15,19,20]. All samples were cleaned as describedabove prior to measurement.

Corrosion tests were conducted at OCP in the 0.2 M NaClelectrolyte and followed in situ by impedance measurements. TheOCP value was recorded for 20 min and then the impedance spectrawere measured for the next 10 min. The procedure was repeated after1 h of immersion and then every hour for a total immersion period of6 h. Frequencies between 100 kHz and 10 MHz were used with anamplitude signal set to 10 mV to guarantee a linear response.

3. Results and discussion

3.1. Surface and interface analysis of pristine coated samples

3.1.1. ToF-SIMS depth profilingFig. 1 shows the ToF-SIMS depth profiles obtained for the pristine

50 nm chromium and tantalum oxide layers grown by FCAD on thecarbon steel alloy substrate. The selected ionswere 12C−, 17OH−, 18O−,28Si−, 35Cl−, 52Cr−, 56Fe−, 64CCr−, 84CrO2

−, 88FeO2−, 91FeCl−, 181Ta−

and 213TaO2−. 18O is the naturally occurring oxygen isotope recorded

since the 16O− signal was close to saturation. The ion intensities arepresented in logarithmic scale in order to emphasize the low intensitysignals, and plotted versus Cs+ sputtering time. Three regions, markedon the profiles, are distinguished. The coating region is defined from thebeginning of the profile until the position where the intensity of ionscharacteristic for the coating (CrO2

− or TaO2−) decreases and that of the

ions characteristic for the substrate (Fe−) increases. It is followed by theinterfacial region and then the substrate region, the transition betweenthese two being set at the position where the intensity profiles of thecoating oxide ions (CrO2

− or TaO2−) intersect the intensity profiles of

ions characteristic of the substrate (Si− and/or Fe−).Starting from the outermost surface, the coating region for the

chromium oxide sample is characterized by a steady intensity of theCrO2

−, Cr−, OH− and 18O− ions (Fig. 1(A)). This shows the growth of afilm with no in-depth variation of the stoichiometry. Some organicand OH contamination is evidenced in this region by the C− and OH−

ion profiles. Peaking at the outermost coating surface because ofexposure to air, the C contamination decreases to a minimum in themiddle of the coating region, showing the penetration of the organicenvironmental contamination in the coating. Afterwards an increaseis observed before reaching the onset of the interfacial region,indicating the influence of the substrate surface residual organiccontamination in the first stage of deposition. The OH profile alsopeaks at the outermost coating surface due to environmental

Fig. 1. ToF-SIMS negative ion depth profiles for the pristine A) 50 nm chromium oxideand B) 50 nm tantalum oxide layers deposited by FCAD on the 100Cr6 carbon steelsubstrate.

3905B. Díaz et al. / Surface & Coatings Technology 206 (2012) 3903–3910

contamination but, in contrast, it does not vary with depth in the bulkcoating region showing that the OH species mostly originates fromwater vapor traces during deposition. Trace Cl contamination,brought by the deposition process (possibly the Cr target), is alsoobserved to not vary with depth.

After about 220 s of sputtering, the onset of the interfacial regionis reached. This region is characterized by well-marked peaks in theintensity profiles of the C− and CCr− ions, showing the presence of acarbidic chromium layer. The presence of this carbidic interface layersuggests interaction of the first deposited metallic layers (Cr particles)with the residual organic contamination left on the substrate surface bythe pre-etching stage of the deposition process. Note that only a shallowpeak is observed at the interface in the profile of the FeO2

− ions, showingthat pre-etching was quite effective in removing most of the nativeoxide at the substrate surface prior to deposition. This is a markeddifference with coatings deposited on the same carbon steel alloy byatomic layer deposition (ALD) for which a substrate surface oxide film(containing Fe and Cr) was systematically observed at the interface[15–18].

The tantalum oxide sample also presents some organic and OHenvironmental contamination at the outermost surface, removedafter ~10 s of sputtering (Fig. 1(B)). Then the TaO2

−, OH− and 18O− ionprofiles remain steady in the bulk of the coating, showing no in-depthvariation of the stoichiometry of the deposited oxide (the observedoscillations are assigned to charging effects caused by sputtering). TheC− ion intensity is stable through the whole coating in contrast to thatof the chromium oxide sample. At the coating surface, it shows a lowerpenetration of environmental contamination and thus a possibly lowerporosity of this coating, as confirmed by electrochemicalmeasurements

presented below. At the coating/substrate interface, it shows no effect ofthe residual substrate surface contamination in the first deposited oxidelayers. Trace Cl contamination is also observed to be steady in depth.

The onset of the interfacial region is reached after about 240 s ofsputtering. The TaC− and C− ion profiles peak at about 260 s,showing the formation of a tantalum carbidic layer. These peaks arepreceded by a Ta− peak, showing that some metallic tantalumparticles remain unreacted between the formation of the carbidicinterface layer and that of the first oxide layers, thus forming a duplexTa/Ta–C interlayer. The peak observed at the interface in the profile ofthe FeO2

− ions is smaller than for the chromium oxide sample,showing an even more effective removal of the substrate native oxideby pre-etching in the deposition process.

Comparing the bulk coating regions for the two samples, onenotices a markedly higher (factor of ~10) intensity of the FeO2

− ionprofile for the chromium oxide sample. This is in line with theelectrochemical data presented below that indicates a more porousstructure of the chromium oxide film. The FeO2

− ions detected in thebulk coating region probably originate from pinholes in the coatingthrough which the primary ions can reach the exposed substratesurface, and whose surface fraction is larger for the chromium oxidesample. Some side effects of the purity of the chromium and tantalumtargets used for deposition cannot be fully excluded.

3.1.2. XPS surface analysisFig. 2 shows the high resolution XP spectra of the Cr2p and Ta4f

core levels for the chromium and tantalum oxide coatings, respec-tively. The fitting parameters, binding energies (BE) and full width athalf maximum (FWHM) of the component peaks, and the extractedatomic fractions are compiled in Table 2, as well as those for the C1sand O1s spectra (not shown).

For the chromium oxide sample the XP Cr2p spectrum is similar tothose of reference Cr2O3 samples reported in the literature, whichindicates that the surface composition of the Cr oxide thin film ismostly Cr2O3 [21–24]. As previously observed, the Cr2p1/2 and Cr2p3/2peaks are broadened and split by satellite peaks and strong multipletinteractions resulting from the coupling of the unpaired 3d valenceelectrons with the 2p core hole, thus making difficult the unambiguousassignment of the binding energies of the components used for peakfitting [24]. Here three sets of 3/2–1/2 doublet components were usedfor peak fitting in agreement with bibliographic data. The Cr2p3/2 andCr2p1/2 doublet splitting was fixed at 9.8 eV [25] and the 2p3/2/2p1/2intensity ratio was set at 2. The main peak (Cr2p3/2B, 57.3% of the globalsignal) is centered at 576.5 eV and corresponds to Cr(III) in Cr2O3

[26–28]. The lower (Cr2p3/2A) and higher (Cr2p3/2C) BE components arecentered at 575.5 and 577.9 eV, respectively. They could be associatedwith the presence of higher and lower oxidation states, respectively[27,29]. The absence of Cr(0) signal at EB=574.4 eV [26] shows that thepresence, if any, of trace amount of metallic chromium in the film, assuggested by ToF-SIMS (Fig. 1(A)) is not detected by XPS (detectionlimit of ~0.5 at.%).

The O1s spectrum shows a main component (O1sA) at 530.2 eVrelated to the oxygen ions and two lower intensity components athigher binding energy. The first peak (O1sB, shifted by +1.5 eV) andthe second one (O1sC, shifted by +2.9 eV) can be assigned toadsorbed hydroxyl groups and carbonates, respectively [30,31]. TheC1s spectrum confirms organic contamination of the outermostsurface, already seen by ToF-SIMS. This organic contaminationconsists mostly of hydrocarbons (peak C1sA at 285 eV, C\C bonds)and some C\O (peak C1sB at 286.5 eV), C_O (peak C1sC at 288.1 eV)and O\C_O (peak C1sD at 288.9 eV) containing species which aretypical of surface organic contaminants. No trace of iron oxide nor ofCl was detected showing that the intensity measured by ToF-SIMS isbelow the detection limit of the spectrometer (b0.5 at.%).

For the tantalum oxide sample the Ta4f spectrumwas fitted with a4f7/2/4f5/2 spin-orbit doublet splitting fixed at 1.9 eV and an intensity

Fig. 2. High resolution XP spectra of the Cr2p and Ta4f core levels for the chromiumoxide and the tantalum oxide layers, respectively.

3906 B. Díaz et al. / Surface & Coatings Technology 206 (2012) 3903–3910

ratio set at 4/3 [32,33]. The Ta4f core level shows four doubletsassociated to four oxidation states of tantalum. The doublet D, at thehighest binding energy, 26/27.9 eV, is the highest in intensity (82.2%of the Ta signal). It is assigned to Ta(V) in Ta2O5 [32–37]. The doubletsC at 24.8/26.7 eV and B at 23.9/25.8 eV have been assigned to loweroxidation states, Ta(IV) and Ta(II), respectively [32,34,37,38]. Theyhave lower intensities, 12.2 and 3.3% of the Ta signal, respectively,showing the minor content of the FCAD film in these lower oxides.The doublet A at 21.7/23.6 eV has the lowest intensity (2% of the Tasignal) and may correspond to even lower oxidation states Ta(I) orTa(0). The O2s at 23 eV is also found in the Ta4f core level region(Fig. 2) [35,39].

Table 2Binding energies (BE), full-width at half-maximum (FWHM) and atomic fraction of thecomponent peaks of the XP C1s, O1s, Cr2p and Ta4f core levels for the pristinechromium and tantalum oxide films grown by FCAD.

Chromium oxide Tantalum oxide

BE(eV)

FWHM(eV)

Atomic fraction(%)

BE(eV)

FWHM(eV)

Atomic fraction(%)

C1sA 285 1.14 25.9 285 1.29 34.4C1sB 286.5 1.14 2.3 286.6 1.29 6.3C1sC 288.1 1.14 1.3 288.2 1.29 1.6C1sD 288.9 1.14 1.7 289.3 1.29 2.4O1sA 530.2 1.53 36.9 530.6 1.4 31.8O1sB 531.7 1.53 9.4 532.3 1.4 10.1O1sC 533.1 1.53 2.7 533.6 1.4 4.4Cr2p3/2A 575.5 0.91 2.0 – – –

Cr2p3/2B 576.5 2.01 11.3 – – –

Cr2p3/2C 577.9 3.16 6.5 – – –

Ta4f7/2A – – – 21.7 1.2 0.2Ta4f7/2B – – – 23.9 1 0.3Ta4f7/2C – – – 24.8 1.2 1.1Ta4f7/2D – – – 26 1.08 7.4

The O1s and C1s spectra are similar to those for the chromiumoxide film (Table 2). The major O1s component is that at the lowestbinding energy (530.6 eV), related to the oxygen ions of the film.Hydroxylation is observed at 532.3 eV as well as surface contamina-tion by carbonates at 533.6 eV, in agreement with the ToF-SIMS data.The fitting of the C1s spectrum reveals that the same species as for thechromium oxide film are present as organic contaminants of theoutermost surface of the tantalum oxide coating.

3.2. Sealing performance

The i–E polarization curves for the uncoated and coated samples(10 and 50 nm chromium or tantalum oxide coatings) measured inthe 0.2 M NaCl aqueous solution are presented in Fig. 3. All thesamples display an active corrosion behavior in the studied range ofpotential, as expected from the low Cr content of the alloy. For allcoated samples, a reduction of current density is observed in both thecathodic and anodic branches, reaching more than one order ofmagnitude for the 50 nm tantalum oxide layer.

The corrosion current density (icorr) and the polarization resis-tance (Rp) were measured at the corrosion potential (Ecorr) usingTafel analysis, and the values are reported in Table 3. Their reductionconfirms the enhanced corrosion resistance for the FCAD coatedsamples. Table 3 also compiles the porosity values (P) which refer tothe uncoated surface fraction as calculated using Eq. (1) [15,19,20]where Rp

0 and Rp are the polarization resistances for the uncoated andcoated samples, respectively.

P ¼ R0p

Rp⋅100% ð1Þ

The basic assumption behind this evaluation is that the electro-chemical activity of the measured sample is that of the uncoatedsubstrate surface, as supported by the similar shape of thepolarization curves for the coated and uncoated specimens (Fig. 3).The accuracy of the absolute values may be affected by changes of theactive surface during the test but the procedure is reliable for samplecomparison since all samples were submitted to the same analysisconditions, including contact time with the electrolyte.

Table 3 shows that growing 10 nm of either of the oxides isinsufficient to obtain appreciable sealing properties. Among thetested thin films, the 50 nm tantalum oxide layer apparently developsa less defective structure thus enabling better sealing of the substratesurface exposed to the aggressive environment. In comparison toother PVD coatings reported in the literature, typically containing Ti–N species and with thickness in the μm scale, the reduction in thepolarization resistance obtained with the 50 nm tantalum oxide layer

Fig. 3. i–E polarization curves of the uncoated and FCAD coated systems measured in0.2 M NaCl solution (dE/dt=1 mV s−1).

Table 3Polarization resistance (Rp) and corrosion potential (Ecorr) values obtained from the i–Epolarization curves. The coating porosity is derived from the Rp values for the uncoatedand coated samples. Pre-etching time is 30 min except where noted.

Ecorr(mV/SCE)

icorr(A cm−2)

Rp

(Ω)Porosity(%)

Bare substrate −753 4.28E−07 1.08E+05 100Cr–O (10 nm) −693 2.08E−07 3.68E+05 29.3Cr–O (50 nm) −698 1.57E−07 4.59E+05 23.5Ta–O (10 nm) −714 1.69E−07 3.48E+05 31.0Ta–O (50 nm) −671 3.48E−08 2.22E+06 4.9Cr–O (10 nm)(90 minute pre-etching)

−480 3.71E−08 1.35E+06 8.0

3907B. Díaz et al. / Surface & Coatings Technology 206 (2012) 3903–3910

grown by FCAD is of the same order of magnitude [10,40,41].Therefore the as-prepared 50 nm Ta–O thin film is most likely denser,with a lower pore density, thus providing equivalent sealingperformance and corrosion resistance at a much lower thickness.

Table 3 also shows that the effect of increasing the thickness onthe sealing property is hardly appreciated for the chromium oxidelayers grown by FCAD. No marked further reduction of thepolarization resistance was obtained when growing the coatingfrom 10 to 50 nm. For tantalum oxide film, the improvement obtainedwith the same thickness variation was more significant since theporosity value was reduced by a factor of ~6. However, in comparisonto oxide layers of the same thickness deposited by ALD theimprovement in porosity obtained by thickening is very small[15–18,42,43]. Indeed, for ALD alumina and tantalum oxide coatingsa reduction in porosity of typically two orders of magnitude (from13.4 to 0.55%) has been obtained with a thickness increment from 10to 50 nm on the same substrate [15,16,18] and 316 L stainless steel[42]. The much lower effect of increasing the thickness for the FCADthin films most likely relates to the well-known columnar growth ofthe PVD layers [39,44], preventing efficient sealing of the microstruc-ture defects formed in the initial stages of deposition.

A possible way to improve the sealing performance of the FCADlayers is better pre-conditioning of the substrate surface prior toactual growth since these layers are very sensitive to surfacecontamination, either oxides or organic pollutants. With moreeffective removal of surface contamination, a more homogeneousand less defective growth in the initial stages of deposition of thecoating can be expected. Fig. 4 compares the ToF-SIMS depth profilesof some selected ions (significant of contamination) for 10 nmchromium oxide layers deposited after 30 and 90 min of pre-etching. The ion intensities and sputtering times cannot be directlycompared to those presented in Fig. 1 since the analysis and sputteredareas were varied in order to increase sensitivity and depth

Fig. 4. ToF-SIMS negative ion depth profiles for 10 nm chromium oxide layers, showingthe effect of the pre-etching time.

resolution. A noteworthy variation is the further (almost complete)removal of the native oxide (FeO2

− ions) in the interfacial region after90 min of pre-etching. A reduction in the hydroxide content (OH− ions)is also observed in the interfacial region. The organic contamination(C− ions) is not significantly affected at the interface but a lowerintensity is measured in the first deposited layers and in the bulkcoating region. This result, coupledwith Table 3 that shows a significantreduction of the porosity with a 90 minute pre-etching time, confirmsthat optimizing the substrate surface cleanliness prior to deposition isinstrumental for promoting the sealing performance of the firstdeposited layers of the coating.

An anodic shift of the corrosion potential for all coated samples isalso observed in Fig. 3 (values are compiled in Table 3). The corrosionpotential of the bare substrate is located at about −0.75 V(SCE), andit is shifted by about 0.05–0.08 V for the 30 minute pre-etched coatedsamples. Ennoblement by modification of the reactive uncoatedsurface is suggested. Most likely, removal of the native oxide presenton the uncoated alloy surface promotes corrosion resistance since theToF-SIMS data barely show the presence of iron oxide at the substratesurface on the coated sample.

3.3. Corrosion behavior

Given their better sealing performance, only the 50 nm thick oxidelayers were submitted to the corrosion test in the neutral 0.2 M NaClaqueous solution. The experiments were carried out in the sameconditions as in a previous work [17] in order to compare the initialcoating stability with that of highly sealing ALD Al2O3 films of thesame thickness deposited on the same carbon steel surface. Longertest would be necessary to assess the durability of the corrosionprotection.

3.3.1. ToF-SIMS depth profiling after immersionThe ToF-SIMS profiles obtained after the corrosion tests are

presented in Fig. 5(A) and (B) for the Cr–O and Ta–O samples,respectively. They are to be compared to profiles of the pristinesamples presented in Fig. 1(A) and (B), respectively. The same criteriawere applied for defining the different regions. No difference of thesputtering time of the coating was observed for any of the studiedsamples, indicating no coating removal and thus no dissolution andgood stability of the FCAD coatings in a neutral Cl-containing solution.This is as expected from the stability range of the Cr2O3 and Ta2O5

oxides and in contrast with that of ALD alumina layers that showed adissolution rate of 7±1 nm h−1 in the same testing conditions whendeposited on the lapped 100Cr6 surface [17].

Some changes were observed by ToF-SIMS in the bulk coatingregions and mostly in the interfacial regions pointing to differentcorrosion resistance behaviors provided by the Cr–O and Ta–Onanocoatings. For both samples, markedly higher intensities of theCl− ion profiles are observed after immersion indicating thepenetration of the electrolyte through the porous coatings until thecoating/substrate interface. The appearance of weak intensities ofcorrosion products containing FeCl− ions in the interfacial region isconsistent with the initiation of localized corrosion (i.e. pitting) at theuncoated substrate sites exposed by the porous coatings.

For the chromium oxide film, the C− ion profile becomes moreintense in the bulk coating region after immersion. This is alsoconsistent with the penetration of organic contaminants accompa-nying the electrolyte entry in the porous coating. The peaks observedin the C− and CCr− ion profiles at the coating/alloy interface of thepristine sample remarkably disappear after the immersion test, mostlikely due to dissolution of the Cr–C carbidic interlayer following thepenetration of the electrolyte via the coating defects. This suggests inthis case a defective growth of the FCAD chromium oxide preferablyover the substrate sites where the Cr–C interlayer was formed prior todeposition of chromium oxide. Some other modifications of the

Fig. 6. Evolution of the impedance spectra (Nyquist plots) for the A) Cr–O and B) Ta–OFCAD coated samples during the immersion test in 0.2 M NaCl at OCP.

Fig. 5. ToF-SIMS negative ion depth profiles for the A) 50 nm chromium oxide andB) 50 nm tantalum oxide samples after immersion in 0.2 M NaCl at OCP for six hours.

3908 B. Díaz et al. / Surface & Coatings Technology 206 (2012) 3903–3910

interfacial region confirm the electrolyte penetration through thecoating defects and the initiation of pitting. A much less sharpdecrease is observed in the intensities of the CrO2

−, Cr−, OH− and18O− ions, consistent with a roughness increase at the coating/substrate interface caused by pitting, and, accordingly, a markedlylonger time (increasing from ~170 to >370 s) needed to sputterthrough the interfacial region.

For the tantalum oxide film, a significant increase of carboncontamination (C− ion profile) is observed only in the external part ofthe coating after immersion, confirming a less porous coatingmicrostructure. The interlayer (Ta/Ta–C duplex in this case) is notmodified after immersion. Thus, and in contrast to the Cr–O film onthe Cr–C interlayer, it is suggested a less defective growth of the Ta–Olayer over the Ta/Ta–C duplex interlayer. No marked differences wereobserved in the shape of the TaO2

−, OH− or 18O− ion profiles, showingno pronounced modification of roughness at the interface. Neverthe-less, an increase (from ~140 up to ~300 s) of sputtering time of thecoating/substrate interfacial region was also observed suggestingsome modification of the interface following the immersion test. Apeak in the intensity of the FeO2

− ion profile can now be observed atthe beginning of the interfacial region, pointing to some accumulationof corrosion products and thus confirming the penetration of theelectrolyte through the coating defects. A slight increase in the OH−

ion intensity in the deeper part of the interfacial region also points tosome accumulation of hydroxyl-containing compounds and confirms,together with for the FeCl− ion profile, the occurrence of localizedcorrosion, however much less marked than on the chromium oxidesample.

3.3.2. EIS analysis during immersionFig. 6 shows the impedance spectra (Bode plots) registered after

0.5, 3 and 6 h in the 0.2 M NaCl solution at OCP. The Ta–O (Fig. 6(B))sample shows a larger impedance than the Cr–O sample (Fig. 6(A)) inagreement to the lower uncoated surface fraction (i.e. lower porosity)measured by polarization and discussed above. For this sample nosignificant change was observed in the spectra during the 6 h of theimmersion test. This confirms the good stability of the sample, inagreement with the minor modifications detected by ToF-SIMS. Forthe Cr–O sample, a slight evolution of the EIS spectra is observedduring the immersion test, also in agreement with the morepronounced modifications observed by ToF-SIMS.

Since the EIS spectra show one time constant in the frequencyrange analyzed, the common Randles equivalent circuit wasemployed for data fitting. The corresponding fitting parameters arereported in Table 4. The charge transfer resistance (Rct), measured atthe low frequency intercept with the x-axis and analogous to thepolarization resistance [45], increases after the first hour of immer-sion for both samples. This initial variation must be related to somemodifications of the substrate surface by the effect of immersion [46].After that the resistance decreases for the Cr–O sample film whereasit continues increasing for the Ta–O sample. The stability of theinterface formed in the first stages of the coating process thus plays akey role in the corrosion performance. As shown by ToF-SIMS, thecarbidic interlayers under the oxide films behave differently dependingon the nature of the FCAD layers. For the Cr–O sample, the initiationperiod preceding corrosion is shorter considering the reduction in thecharge transfer resistance measured after 6 h of immersion, whichpossibly relates to the dissolution of the Cr–C interfacial layer. Incontrast, for the Ta–O sample, the higher chemical stability of the

Table 4EIS fitting parameters for the spectra in Fig. 6 using the Randles equivalent circuit.

OCP(mV)

Re(Ω)

Rct

(Ω)CPE(Ω−1 sn)

n

Bare substrate −672 91.9 3.31E+04 2.96E−05 0.79Cr–O 30 min −684 71.4 1.48E+05 1.15E−05 0.81Cr–O 1 h −722 71.1 1.56E+05 1.16E−05 0.81Cr–O 6 h −730 67.6 1.32E+05 1.48E−05 0.82Ta–O 30 min −630 79.5 2.83E+06 5.35E−06 0.96Ta–O 1 h −652 79.7 3.10E+06 5.43E−06 0.95Ta–O 6 h −668 78.9 3.26E+06 5.63E−06 0.95

3909B. Díaz et al. / Surface & Coatings Technology 206 (2012) 3903–3910

interface including the Ta/Ta–C interlayer extends the initiation perioduntil at least 6 h of immersion.

The capacitance was modeled using a Constant Phase Element(CPE) in order to take into account the deviation from ideal dielectricbehavior due to surface heterogeneities such as roughness or inhomo-geneous distribution of the electrode surface properties [47–50]. Theexact meaning of the CPE parameter cannot be precisely defined. Themeasured capacitances reported in Table 4 show for both coatings aslight increase with immersion time. Since the coating dissolution canbe excluded on the basis of the ToF-SIMS data, this increase must beassociated to a change at the uncoated sites generated by the corrosivesolution. The larger capacitance increase observed for the Cr–O sample(29% vs. 5% for the Ta–O sample) is then in agreement with the morepronounced interface modification shown by ToF-SIMS.

In our previous work on ALD coatings [15], it was established thatfor samples with a high coating porosity, i.e. a large uncoated surfacefraction, the most significant contribution to the capacitance wasgiven by the double layer capacitance, whereas, in the case of a lowporosity, the main contribution was that of the coating capacitance.According to this, for the present coatings which have a significantlyhigher porosity than the ALD alumina layers previously discussed, themeasured capacitance is expected to be dominated by the doublelayer capacitance. Assuming this, the coating porosity can also becalculated using the capacitance values obtained for the uncoated andcoated samples after 30 min of immersion and reported in Table 4.Using this procedure, the obtained porosity values would be ~39 and~18% for the Cr–O and Ta–O samples, respectively. These values, inparticular for the Ta–O sample, are markedly higher than those(Table 3) deduced from the polarization resistance values whereasthe values, ~23 and ~1%, respectively, deduced from the chargetransfer resistance are in reasonably good agreement.

This discrepancy indicates that the double layer capacitance isoverestimated and that the contribution of the coating capacitance inthe global capacitance should also be taken into account. The coatingcapacitance can be theoretically calculated considering the definitionof a flat parallel capacitor as described in Eq. (2).

Ccoat ¼εε0Sd

ð2Þ

where ε refers to the dielectric constant of the layers, typically 12 and27 for bulk Cr2O3 and Ta2O5, respectively [51,52], ε0 corresponds tothe vacuum dielectric constant (8.85·10−14 F cm−1), d is the coatingthickness and S is the coated surface. The theoretical coatingcapacitances are then 7.3·10−8 F and 2.1·10−7 F for the 50 nmchromium and tantalum oxide films, respectively. These values aretoo low to significantly affect the global capacitance since themeasured values in Table 4 are one to two orders of magnitudehigher. It follows that the real coating capacitances for these FCADlayers must be higher than those for bulk oxides. Considering Eq. (2),a deviation of one to two orders of magnitude of the thickness can beexcluded since the actual thickness of the FCAD layers was controlledafter deposition. The remaining possibility to explain the deviation

would be that the dielectric constants of the films are significantlylarger than those commonly referred to for these oxides. Thissuggests that the FCAD oxide nanocoatings studied in this work arepoor dielectrics, suggesting poor insulating properties in line with adefective and porous nanostructure of the coatings indirectlyevidenced by the measurement of the sealing performance.

4. Conclusions

Thin (10 and 50 nm) chromium and tantalum oxide coatings grownby FCAD for corrosion protection of carbon steel have been studiedusing ToF-SIMS cross-section and XPS surface analysis, and voltamme-try and EIS electrochemical analysis. Pre-etching by ion bombardmentprior to actual oxide growth in the deposition process allowed us toreduce the native oxide layer formed on the steel surface to traces ofiron oxide. The presence of a carbidic interlayer at the interfacewith thecarbon steel substrate is evidenced by ToF-SIMS. It results from thereaction between the first deposited metallic particles and the residualcarbon contamination of the substrate surface left by pre-etching. Thebulk coatings mostly consist of stoichiometric oxides, Cr2O3 or Ta2O5,with no significant in-depth variation. Surface and bulk contaminationby hydroxyl and organic groups was observed for both coatings.

Electrochemical measurements evidenced a relatively large coatingporosity or uncoated surface fraction for the 50 nmfilms, decreased by afactor of ~5 for the better sealing tantalum oxide coatings. Increasingthe film thickness from 10 to 50 nm is not as effective to improve thesealing properties as for ALDoxidefilms,whichhas been assigned to theknown columnar growth of the PVD films preventing good sealing atgrain boundaries. The duplex structure (Ta/Ta–C) of the carbidicinterlayer is thought to promote a less defective growth of tantalumoxide than the single Cr–C interlayer for chromium oxide, therebyimproving the sealing properties. The dielectric constants found for the50 nmCr2O3 and Ta2O5 coatings are higher than the common values forthese oxides, suggesting poor insulating properties in line with adefective and porous nanostructure of the coatings.

An excellent stability (i.e. no dissolution) was observed for bothoxide nanocoatings in neutral 0.2 M NaCl at OCP, as opposed to thedissolution of 50 nm ALD alumina grown on the same carbon steelsubstrate. Penetration of the electrolyte to the interface with thecarbon steel surface caused the dissolution of the Cr–C interlayer butnot that of the Ta/Ta–C interlayer. As a result an improved corrosionresistance was measured for the sample covered with the tantalumoxide coating.

Acknowledgments

The research leading to these results has received fundingfrom the European Community's Seventh Framework Programme(FP7/2007–2013) under grant agreement no. CP-FP 213996-1(CORRAL). Region Ile-de-France is acknowledged for partial supportfor the ToF-SIMS equipment.

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