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Electrochimica Acta 56 (2011) 10516– 10523

Contents lists available at ScienceDirect

Electrochimica Acta

j ourna l ho me p ag e: www.elsev ier .com/ locate /e lec tac ta

lectrochemical and time-of-flight secondary ion mass spectrometry analysis ofltra-thin metal oxide (Al2O3 and Ta2O5) coatings deposited by atomic layereposition on stainless steel

elén Díaza, Jolanta Swiatowskaa, Vincent Mauricea,∗, Antoine Seyeuxa, Bernard Normandb,mma Härkönenc, Mikko Ritalac, Philippe Marcusa

Laboratoire de Physico-Chimie des Surfaces, CNRS (UMR 7045) - ENSCP (Chimie-ParisTech), 11 rue Pierre et Marie Curie, F-75005 Paris, FranceUniversité de Lyon, INSA de Lyon, MATEIS (UMR CNRS 5510), Bât. L. de Vinci, F-69621 Villeurbanne, FranceLaboratory of Inorganic Chemistry, University of Helsinki, P.O. Box 55, FIN-00014 Helsinki, Finland

r t i c l e i n f o

rticle history:eceived 15 November 2010eceived in revised form 17 February 2011ccepted 18 February 2011vailable online 26 February 2011

eywords:xide coatingsorrosion protectionLD

a b s t r a c t

Ultra-thin (5–50 nm) layers of aluminium and tantalum oxides deposited by atomic layer deposition (ALD)on a stainless steel substrate (316L) for corrosion protection have been investigated by electrochemicalmethods (linear scan voltammetry, LSV, and electrochemical impedance spectroscopy, EIS) and time-of-flight secondary ion mass spectrometry, ToF-SIMS. The effects of the deposition temperature (250 ◦Cand 160 ◦C) and coating thickness were addressed. ToF-SIMS elemental depth profiling shows a markedeffect of the organic and water precursors used for deposition and of the substrate surface contaminationon the level of C and OH trace contamination in the coating, and a beneficial effect of increasing thedeposition temperature. The polarization data show a decrease of the current density by up to fourorders of magnitude with increasing coating thickness from 5 to 50 nm. The 50 nm films block the pitting

ISoF-SIMS

corrosion in 0.8 M NaCl. The uncoated surface fraction (quantified from the current density and allowing aranking of the efficiency of the coating, also confirmed by the capacitance and resistance values extractedfrom the EIS data) was 0.03% with a 50 nm thick Al2O3 film deposited at 250 ◦C. The correlation betweenthe porosity values of the coatings and the level of C and OH traces observed by ToF-SIMS points to amarked effect of the coating contaminants on the sealing performance of the coatings and on the corrosionresistance of the coated systems.

. Introduction

Metallic surfaces can be protected against corrosion by coatingshich separate the metal or alloy from the aggressive environment.

he interest for the application of thin films as protective coatingsas increased recently. Several techniques have been applied for theeposition of thin layers, including magnetron sputtering, chemicalnd physical vapor deposition methods [1–5]. However, the pres-nce of defects or pinholes in the coatings cannot be completelyvoided and limits the corrosion protection performance [3,6].

Atomic layer deposition (ALD), a thin film deposition methoderived from chemical vapor deposition, is an ideal candidate torow advanced ultra-thin and non porous coatings for corrosion

rotection. ALD allows the growth of high-quality, thin and uni-orm films with excellent conformality [7–9]. The film uniformityxtends not only over flat substrate surfaces but also into very nar-

∗ Corresponding author.E-mail address: vincent-maurice@chimie-paristech.fr (V. Maurice).

013-4686/$ – see front matter © 2011 Elsevier Ltd. All rights reserved.oi:10.1016/j.electacta.2011.02.074

© 2011 Elsevier Ltd. All rights reserved.

row holes and trenches. When the deposition parameters such astemperature, precursor pulse and purge times are properly cho-sen, each reaction step reaches saturation, making the growthself-limiting and the amount of material deposited in each stepuniform over the surface. The saturative ALD reactions also ensurereproducible and precise thickness control simply by counting thenumber of reaction cycles. Commercially ALD is used mostly inmicroelectronics but also for making invisibly thin antitarnish-ing coatings on silver jewellery [10]. Despite the excellent filmproperties, studies on using ALD for depositing corrosion protec-tion coatings have remained quite limited. The very first study[7] showed that a few hundred nanometer thick ALD aluminium,titanium and tantalum oxide coatings give excellent protectionto stainless steel against corrosion. Ultra thin, about 50 nm thick,TiO2 layers deposited by ALD, have also been proved to efficientlyimprove the corrosion performance of stainless steel [11].

In this paper the corrosion resistance properties of ultra-thin(from 5 to 50 nm) amorphous aluminium and tantalum oxide coat-ings prepared by ALD are presented and discussed. The influenceof the deposition temperature (250 ◦C vs. 160 ◦C) is addressed.

B. Díaz et al. / Electrochimica Acta 56 (2011) 10516– 10523 10517

Table 1Chemical composition (wt%) of the 316L stainless steel.

Tb(pcfc

2

2

pwfimww

2

Ad(tp20w5i

fsmaT

TP(o

Element C Cr Ni Mo Mn

wt% Max. 0.03 18–20 13 2.7 Max. 2

he elemental depth distributions in the coatings were analyzedy means of time of flight secondary ion mass spectroscopyToF-SIMS). Potentiodynamic polarization measurements wereerformed to analyze the corrosion protection and to quantify theoating sealing performance, in particular the uncoated surfaceraction. Electrochemical impedance spectroscopy was applied toharacterize the coating quality.

. Experimental

.1. Material and surface preparation

The chemical composition of the 316 stainless steel substrates isresented in Table 1. The sample surface was mechanically polishedith successive 9, 3 and 1 �m diamond suspension followed by anal polishing with a SiO2 suspension (0.04 �m). The roughnesseasured by profilometry was 1.7 nm. Before coating, the samplesere cleaned in an ultrasonic bath of isopropanol for 5 min, rinsedith ethanol and dried by blowing with compressed air.

.2. Coating preparation by atomic layer deposition

The coatings were prepared with a Picosun SUNALE R-150LD reactor. The precursors employed for the deposition proce-ure were tantalum ethoxide (Ta(OC2H5)5), trimethylaluminiumAl(CH3)3) and H2O. Ta(OC2H5)5 was evaporated at 140 ◦C whilehe other two precursors were kept at room temperature. The tem-eratures in the deposition chamber were selected to be 160 and50 ◦C. The deposition rates were about 0.048 nm/cycle and about.1 nm/cycle for Ta2O5 and Al2O3, respectively. The ALD reactoras operated with a constant nitrogen flow at a pressure of about

mbar. The nitrogen carrier and purge gas flow was 600 sccm withmpurity content below 10−5 (impurity flow below 0.006 sccm).

The thickness of the coatings was measured from the sur-ace of a silicon wafer coated simultaneously with the steel

ubstrates. The thickness of the thicker coatings (≥50 nm) waseasured with an UV-VIS spectrophotometer (Hitachi U-2000)

nd simulated using the Ylilammi and Ranta-aho [12] software.he thickness of the thinner samples was measured with X-ray

able 2assive current density (jpass) and porosity (PLSV) values for the uncoated and coated samaccording to Eq. (1)). The EIS fitting parameters obtained using the equivalent circuit presf the CPE values obtained by fitting. The Ccoat values obtained after subtracting the Cnativ

Sample LSV

jpass (A cm−2) PLSV (%)

Un-coated 8.6E−07 –

Al2O3 160 ◦C 5 nm 3.0E−07 35.48

10 nm 2.7E−07 31.45

20 nm 4.8E−09 0.5650 nm 4.2E−09 0.49

Al2O3 250 ◦C 5 nm 1.3E−07 15

10 nm NAa NAa

20 nm 6.9E−10 0.08

50 nm 2.4E−10 0.03

Ta2O5 250 ◦C 5 nm 1.2E−07 13.80

10 nm 2.2E−07 26.05

20 nm 2.0E−07 23.01

50 nm 1.2E−09 0.14

a NA – procedure not applicable.

Si P S N Fe

Max. 0.75 Max. 0.04 Max. 0.03 Max. 0.1 Bal.

reflectance (XRR, Bruker AXS D8 Advance) and modelled usingLeptos 7.03.

Table 2 shows the thicknesses and deposition temperatures ofthe prepared samples.

2.3. ToF-SIMS depth profiling

Elemental depth profiles were obtained using a ToF-SIMS 5spectrometer (IonToF). The spectrometer was run at an oper-ating pressure of 10−9 mbar. A pulsed 25 keV Bi+ primary ionsource was employed for analysis, delivering 1.8 pA of currentover a 175 �m × 175 �m area. Depth profiling was performed witha 2 keV Cs+ sputter beam giving a 100 nA target current overa 400 �m × 400 �m area. Data acquisition and post-processinganalyses were performed using the Ion-Spec software. Negativeion profiles were recorded because they are more sensitive tofragments originating from oxide matrices. The samples wereexamined in exactly the same sputtering and analysis conditions,allowing a direct comparison of the intensities of the profiles.The ToF-SIMS analyses were performed on pristine samples notexposed to the corroding electrolyte. Before the measurementsthey were cleaned in ethanol as described below.

2.4. Electrochemical measurements

The electrochemical measurements were carried out in a con-ventional three-electrode cell using a saturated calomel electrode(SCE) and a platinum wire as reference and counter electrodes,respectively. The working electrode area was 0.44 cm2. All theexperiments were performed at room temperature in a 0.8 M NaClsolution prepared with ultra pure Millipore® water (resistivity>18 M� cm) and reagent grade chemicals (NaCl Analar Norma-pur analytical reagent VWR® BDH Prolabo®). The electrolyte wasdeaerated by bubbling argon gas for 30 min prior to the measure-ments. The pH was 7.2 ± 0.2 before the corrosion tests. An Autolab

PGSTAT30 was used.

The samples were cleaned with ethanol in an ultrasonic bathfor 10 min and dried with compressed air before the electrochem-ical measurements. Before starting the DC and AC electrochemical

ples calculated from the linear scan voltammetry measurements in 0.8 M Naclaq

ented in Fig. 6 are also compiled. The capacitance values, Eq. (4), are shown insteade are also given.

EIS

Re (�) Rc (�) Cc (F) N Ccoat (F)

22.2 7.8E+06 6.3E−06 0.95 –

21.1 7.2E+06 3.8E−07 0.92 4.1E−0718.3 1.1E+07 2.6E−07 0.91 2.7E−0720.6 1.4E+08 1.1E−07 0.97 1.1E−0723.5 3.6E+08 8.9E−08 0.97 8.1E−08

31.2 1.4E+07 7.4E−07 0.95 8.4E−0722.9 1.7E+09 3.3E−07 0.99 3.5E−0722.2 2.8E+09 1.7E−07 0.99 1.8E−0721.5 6.9E+09 7.0E−08 0.99 7.0E−08

23.2 3.8E+06 2.4E−06 0.92 3.9E−0631.6 4.2E+06 3.2E−06 0.88 6.4E−06NA NA NA NA NA20.8 8.2E+07 1.6E−07 0.98 1.7E−07

1 ica Acta 56 (2011) 10516– 10523

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104

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(cou

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OHO

Al

AlCCrFe

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AlO2

CrO2

FeO2

NiO2

coa tin g

sub str ate

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250ºC

B

0 100 20 0 300 400 50 0

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Spu ttering time / s

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(cou

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coating

substrateinter face

C

OHO

Al

AlC

CrFe

Ni

AlO2

CrO2

FeO2

NiO2

50 nm Al2O3

160 ºC

Fig. 1. ToF-SIMS negative ion depth profiles for the 50 nm Al2O3 coatings preparedat (A) 250 ◦C and (B) 160 ◦C.

0 10 20 30 40 5010

10

50 nm Al2O

3

250ºC

Inte

nsity

(cou

nts)

OH pro file

50 nm Al2O

3

160ºC

0518 B. Díaz et al. / Electrochim

xperiments, the open circuit potential (OCP) was measuredn order to reach a stable potential. Then, the electrochemicalmpedance spectroscopy (EIS) measurements were performed atCP, in a frequency range between 105 and 10−2 Hz, with the ampli-

ude of the exciting signal set at 10 mV. The linear scan voltammetryLSV) tests were performed after the EIS measurements with a scanate of 1 mV s−1 from −0.9 V up to 0.5 V.

.5. Porosity evaluation

The sealing performance of the coating was quantified by eval-ating the surface fraction of the uncoated substrate exposed tohe electrolyte. It is hereafter defined as the coating porosity.efects, pores and/or pinholes, present in the coating are assumed

o connect the bulk electrolyte to the substrate surface wherehe corrosion reactions proceed. Several electrochemical methodsased on different estimations have been previously proposed tovaluate the coating porosity. They are based on a comparison ofhe polarization resistances obtained from impedance data [13,14]r from potentiodynamic tests [1,2,4] of uncoated and coated sam-les. A comparison of the current density values for a fixed potentialalue has also been proposed by Landolt et al. for passive systems15].

In this study, the evaluation of porosity was based on the linearcan voltammetry measurements. For the passive systems pre-ented in this work, the porosity is calculated as a ratio of theassive current density ipass for the coated sample to the passiveurrent density i0pass of the uncoated sample (Eq. (1)).

= ipass

i0pass× 100% (1)

he passive current was always taken at a potential 500 mV morenodic than the corrosion potential.

. Results and discussion

.1. Time-of-flight secondary ion mass spectrometry

Fig. 1 shows the ToF-SIMS depth profiles for the 50 nm thickl2O3 coatings prepared at 250 and 160 ◦C. The distribution of ions

s plotted versus the sputtering time with Cs+ ions. The intensity islotted in logarithmic scale in order to emphasize the low intensityignals.

Three regions can be identified: (i) the coating, (ii) the coat-ng/substrate interface, and (iii) the alloy substrate. The coatingegion is characterized by a plateau of the O−, Al− and AlO2

−

on profiles. At about 330 s, the intensity of the Fe, Cr and Ni-ontaining ions starts to increase marking the beginning of theoating/substrate interfacial region in which the intensity of the Al-ontaining ions drops. The intensity of the Fe, Cr and Ni-containingons becomes stable after approximately 450 s of sputtering, indi-ating the end of the interfacial region and entry in the bulkubstrate. For both samples the width of the interfacial region is20–130 s. This similar width is expected for coatings with similarorphology, both being amorphous and conformal, and deposited

n substrates having a similar roughness. The substrate roughness,y causing shadowing effects of sputtering, enlarges the width ofhe interfacial region.

In the coating region, the intensity of the O−, Al− and AlO2− ion

rofiles recorded with the same primary ion beam current is theame for the two samples indicating a similar stoichiometry of the

lumina coatings prepared at 250 and 160 ◦C. However, the OH−

on profiles at the extreme and near surface region, presented inagnification in Fig. 2, show a higher OH contamination for the

oating prepared at lower deposition temperature. The OH content

Sputte ring t ime / s

Fig. 2. ToF-SIMS OH− ion depth profiles of the extreme and near surface regions forthe 50 nm Al2O3 coatings.

ca Acta 56 (2011) 10516– 10523 10519

idbtl

oi1Trftlltawofep

fisbtoiaip

ipLis

tttciadatmt

TctstiAIyracict

0 50 10 0 150 20 0 25 0 30 0 35 0

10

10

10

10

10

Inte

nsity

(cou

nts)

Sputtering ti me / s

coat ingsubstrateinterface

C

OH

O

TaC

TaO2

Ni

Cr

Fe

FeO2

NiO2

CrO2

50 nm Ta2O5

250 ºC

B. Díaz et al. / Electrochimi

n the bulk of the coating is higher, by a factor of 2, in the coatingeposited at 160 ◦C. This is an expected result as it is known thatelow about 200 ◦C the impurity contents in ALD oxide films starto increase quite strongly with decreasing temperature because ofess complete reactions [9,16].

Fig. 2 also presents the OH contamination at the extreme surfacef the coating. The sputtering time needed to reach a stable OH− ionntensity (i.e. the bulk of the coating) is longer after deposition at60 ◦C (around 20 s) than after deposition at 250 ◦C (around 12 s).he presence of OH-containing species at the extreme surface iselated to the exposure of the sample to ambient air that causes sur-ace hydroxylation and hydration of the alumina film. The longerime needed to reach the OH bulk level in the aluminium oxideayer deposited at lower temperature shows that surface hydroxy-ation/hydration is higher and that the OH ions penetrate deeper inhe sub-surface region of this Al2O3 layer. This difference observedt the coating surface can be related to the coating porosity. Indeed,ith increasing porosity, one can expect that more defects (nano-

r micro-channels) are present in the coating and thus promote sur-ace hydroxylation/hydration and the penetration of OH ions whenxposed to ambient conditions. As a result of this analysis, a higherorosity is expected for the Al2O3 coating prepared at 160 ◦C.

Carbon contamination is also shown by the C− and AlC− ion pro-les. It is observed at the extreme surface but decreases after a feweconds of sputtering, thereby implying its origin being hydrocar-on contamination from ambient. The C− and AlC− ion profiles inhe bulk coating indicate the presence of some unreacted fragmentsf the metal organic precursor Al(CH3)3 trapped inside the coat-ng. This carbon contamination decreases with the deposition timend is reduced at higher deposition temperature, showing that thencrease of temperature promotes elimination of residues of bothrecursors used for deposition of the Al2O3 coatings.

A higher amount of carbon contamination is observed in thenterfacial region. This is assigned to the organic contaminationresent on the substrate surface before the coating deposition.ikewise, the OH contamination increases in the interfacial regionndicating hydroxylation of the stainless steel surface prior to depo-ition.

The Fe, Cr and Ni-containing ion profiles all display a peak inhe interfacial region assigned to the presence of a native oxide onhe substrate surface before deposition. The growth of a spurioushermal oxide in the initial stages of the deposition, prior toomplete sealing of the substrate surface can also contribute to thenterfacial oxide. The slightly smaller width of the FeO2

−, CrO2−

nd NiO2− peaks observed for the sample prepared at the higher

eposition temperature suggests the formation of a thinner oxide,nd thus faster sealing of the substrate surface at higher depositionemperature. The deposition temperature does not markedly

odify the composition of the interfacial region as indicated byhe relative intensities of the FeO2

−, CrO2− and NiO2

− peaks.Fig. 3 presents the ToF-SIMS depth profile for the 50 nm thick

a2O5 coating. The profile intensities and width cannot be directlyompared to those obtained for the Al2O3 coating of the samehickness since the oxide matrices are different, which affects theputter yields of all ions. However, marked differences of intensi-ies (larger than a factor of 10) are observed for the C− and OH−

on profiles in the bulk coating region compared to those of thel2O3 coating prepared at the same temperature (50 nm, 250 ◦C).

t has been previously measured by elastic recoil detection anal-sis that the carbon and hydrogen contents were <0.2 and 1 at.%,espectively, in the Al2O3 film deposited at 250 ◦C from Al(CH3)3nd water, and 0.3–1.1 and 0.6–1.5 at.%, respectively, in the Ta2O5

oating deposited at 300 ◦C from Ta(OC2H5)5 and water [16]. Thisndicates that higher amounts of unreacted fragments of the pre-ursors are retained inside the oxide matrix during the growth ofhe Ta2O5 coating. Similarly to the Al2O3 coating, the interfacial

Fig. 3. ToF-SIMS negative ion depth profile for the 50 nm Ta2O5 coating prepared at250 ◦C.

coating/substrate region is contaminated by a surface native oxideand carbon species.

3.2. Linear scan voltammetry

Fig. 4 shows the potentiodynamic curves for the bare andcoated samples. For the bare 316L alloy, a peak assigned to theactive/passive transition is observed at around −0.2 V, in verygood agreement with previous measurements [17]. Its presence isassigned to the reduction (most likely partial) of the native oxideoccurring during cathodic polarization. For all the coated samples,this reaction is hindered. At the pitting potential of 0.4 V, the sharpincrease in the current marks the initiation and growth of pits forthe uncoated sample. For the samples covered by the thinner (5 and10 nm) Al2O3 coatings deposited at 250 ◦C (Fig. 4A), current fluctu-ations related to metastable pitting were observed in the passiverange and preceded stable pitting. However, for the thicker (20 and50 nm) coatings deposited at 250 ◦C, no metastable nor stable pit-ting could be observed indicating efficient protection and blockingof localized corrosion in the explored range of potential. For theAl2O3 coatings prepared at 160 ◦C (Fig. 4B), metastable pitting andstable pitting were also observed for the 5 and 10 nm thick films,indicating a less efficient corrosion protection. Blocking of localizedcorrosion was achieved only with the 20 nm thick coating.

For both deposition temperatures, the curves in Fig. 4 showthe decrease of the anodic current density with increasing coat-ing thickness, showing that the coating not only blocks localizedcorrosion but also markedly decreases the corrosion rate in thepassive state. Table 2 compiles the passive current density val-ues as well as the porosity values calculated from Eq. (1). Due tothe noise observed during the measurement of the 10 nm Al2O3film (Fig. 4A), passive current density and porosity values could notbe extracted for this sample. As expected, at both deposition tem-peratures, the lowest porosity values are obtained for the thickestcoatings prepared. The results presented in Table 2 show a sig-nificant improvement of the coating sealing performance (lowerporosity) also with the increase of the deposition temperature. Abetter sealing performance (porosity of 0.03%) is obtained with the50 nm thick coating prepared at 250 ◦C than with the 50 nm thick

coating prepared at 160 ◦C (porosity of 0.5%). The porosity orderobtained for the two 50 nm thick coatings prepared at 250 and160 ◦C is in agreement with the ToF-SIMS data that show a deeperpenetration of hydroxide anions assigned to a more defective struc-

10520 B. Díaz et al. / Electrochimica Act

A

-1.2 -1. 0 -0.8 -0.6 -0.4 -0.2 0. 0 0.2 0.4 0.6

1E-1 1

1E-1 0

1E-9

1E-8

1E-7

1E-6

1E-5

1E-4

1E-3

ba re subst rate 31 6L Al O 5nm 250º C Al O 10nm 250º C Al O 20nm 250º C Al O 50nm 250º C

j / A

cm

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ba re subs trate 316 L Al O 5nm 160 ºC Al O 10n m 160º C Al O 20n m 160º C Al O 50n m 160º C

C

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1E-11

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j / A

cm

-2

E vs SCE / V

bare substrate 316 L Ta2O5 5nm 25 0ºC Ta2O5 10nm 250ºC Ta2O5 20nm 250ºC Ta2O5 50nm 250ºC

F(

tTth

The equivalent circuit shown in Fig. 6A is presented as the mostreasonable for the fit of the EIS data, either for the uncoated or the

ig. 4. Polarization curves of the uncoated and Al2O3 coated samples prepared atA) 250 ◦C and (B) 160 ◦C and for the Ta2O5 coated samples prepared at (C) 250 ◦C.

ure of the lower temperature coating. The comparison with theoF-SIMS data also suggests that the more defective structure of

he coating prepared at lower temperature may be related to theigher amount of precursor fragments contaminating the coating.

a 56 (2011) 10516– 10523

Fig. 4C shows the potentiodynamic curves obtained with theTa2O5 coatings. A significant improvement of corrosion resis-tance is observed only with the 50 nm thick coating that blocksmetastable and stable pitting and markedly reduces the corrosionrate in the passive state. For the thinner 20 nm coating, metastablepitting is again observed in the anodic range but no stable pittingoccurs in the tested potential range. This fact suggests improvedrepassivation and higher resistance to stable pitting with the Ta2O5layers than with the Al2O3 layers. Still the higher current denotesa lower sealing performance than the 20 nm thick Al2O3 coatingdeposited at 250 ◦C. Table 2 allows quantitative comparison of thecoating sealing performance, confirming a drastic difference (fac-tor of ∼300) between the 20 nm thick Ta2O5 and Al2O3 coatingsprepared at 250 ◦C. For the 5 nm thick coatings, this difference isreduced to a factor of 5. Again the comparison with the ToF-SIMSdata points to a key role of the C-containing precursor fragmentscontaminating the coating and in higher concentration in the Ta2O5than in the Al2O3 coating as shown above. The decreasing differenceof the sealing performance between the Ta2O5 and Al2O3 coat-ings with increasing thickness is then explained by the decreasingcontamination by precursor fragments, as shown by the ToF-SIMSdepth profile analysis. It may also indicate lower nucleation density,and thereby later closure, of Ta2O5 compared to Al2O3.

3.3. Electrochemical impedance spectroscopy

EIS has been widely used in the last two decades to character-ize the corrosion behavior of coated metallic surfaces, and severalattempts to characterize the coatings quality by EIS have been made[18–21]. Fitting with an adequate equivalent electric circuit is nec-essary for the interpretation of the impedance spectra.

Fig. 5A and B shows the impedance data for the Al2O3 coatingspresented as Bode plots (impedance module and phase angle vs.log(freq)). Only one time constant can be distinguished in the spec-tra for the uncoated as well as for the coated samples. The phaseangle is close to 90◦ which shows a strong capacitive responseeither from the native oxide present on the uncoated samples orfrom the coated samples [22]. The capacitive response resultingfrom the native oxide can be considered as constant for all analyzedsamples if we assume no thickness variation of the native oxideduring coating. However, the data show for both deposition tem-peratures an increasing capacitive behavior at higher frequenciesas the thickness of the coating increases, which is directly related tothe dielectric properties of the Al2O3 coating. As the coating thick-ness increases, the corresponding dielectric capacitance decreasesfollowing the description for a parallel plate capacitor (Eq. (2)):

C = εε0S

d(2)

where S is the surface area covered by the coating (cm2), d thecoating thickness (cm), ε the dielectric constant for the mate-rial inside the capacitor (ε = 9 for Al2O3 [23] and ε = 25 for Ta2O5[24]), and ε0 corresponds to the vacuum dielectric constant(8.85 × 10−14 F cm−1). So the time constant, �, related to this capac-itance and defined as (RC)−1 where R is the coating resistance willbe shifted to higher frequencies with increasing coating thickness.

An equivalent circuit with two time constants representing thecoating and the substrate is frequently employed in most of thepapers when characterizing the quality of the coating [1,4,19,25].However, without a clear indication of the existence of two timeconstants as shown by the Bode plots presented in Fig. 5, it is notpossible to clearly distinguish the coating and the substrate [18,22].

coated samples. Native oxide layers formed on the stainless steelsurface are highly stable [26–28] with high capacitive responses

B. Díaz et al. / Electrochimica Acta 56 (2011) 10516– 10523 10521

A

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bare su bstrate 316L Ta O 5nm 250 ºC Ta O 10nm 250 ºC Ta O 20nm 250 ºC Ta O 50nm 250 ºC

Fig. 5. Bode plots for the uncoated and Al2O3 coated samples prepared at (A) 250 ◦Ca

dsctaatp

nd (B) 160 ◦C and for the Ta2O5 coated samples prepared at (C) 250 ◦C.

ue to their low thicknesses (around 2 nm [29–31]). They are con-idered as inert in our EIS measurements performed at OCP withoutathodic pre-treatment. Thus, no elements have been introduced inhe equivalent circuit to account for active corrosion. Moreover the

nalyzed frequencies (ω) were not low enough to get informationbout the corrosion process and only information about the dielec-ric behavior of the coating can be obtained [22]. Therefore theroposed equivalent circuit presented in Fig. 6 is the most appro-

Fig. 6. (A) Equivalent electric model employed for fitting the impedance spectra (seetext for details). (B) Diagram describing the ALD coating on the metallic substrate.

priate and the information concerning the coating as a dielectricmaterial can be extracted [23,32].

For the proposed electric model, Re, the limit at high frequency,describes the electrolyte resistance, Rc refers to the electric resis-tance of the layer (oxide coating + native oxide on the substrate) andthe Cc-CPE corresponds to the capacitance not only of the dielec-tric film deposited on the sample but also of the native oxide onthe substrate [22]. For the thicker oxide coatings with lower dielec-tric capacitance, the contribution of the native oxide, with higherdielectric capacitance due to the lower thickness, can be neglectedsince both capacitances are in series. In this case, only the lowercapacitance will be considered.

Constant phase elements (CPEs) are commonly used in data fit-ting to describe the non-ideal capacitive behavior of the system.The CPE has no clear physical meaning but it allows for an evalua-tion of the complexity of the studied system. The impedance for aCPE is defined in Eq. (3) [33]:

ZCPE = 1Q (jω)n (3)

where Q is a constant value, independent of the frequency, obtaineddirectly from the fitting, and the factor n is the CPE power, whichis related to the angle ((1 − n) × 90) that evaluates the Nyquistplot depression below the x-axis. The factor n is related to thesurface heterogeneity such as irregularity, roughness or porosity[21,34]. For n = 1 the CPE describes an ideal capacitor and for n closeto 0.5 the CPE represents a Warburg impedance with a diffusioncharacter. The smoother and cleaner the surface, the closer is theparameter n to unity [33,35]. The Q and n parameters are obtainedfrom the fitting and the real equivalent capacitance value can beeasily calculated from the Brug equation [36] (Eq. (4)).

C = Q 1/nRe(1−n)/n (4)

Some other equations have been proposed to obtain the realcapacitance from the Q value [37] but Huang et al. [38] have proventhat the most realistic results were obtained following the Brugformula [36].

The parameters of the EIS data fit for the Al2O3 coated samplesare compiled in Table 2. The capacitance values for the coatings

(Cc following Eq. (4)) are presented instead of the CPE-Q obtaineddirectly from the fitting. The good correlation between the theo-retical values of the capacitance for the oxide coatings (calculatedfrom Eq. (2), and using the coating coverage, 1 − P, deduced from the

10522 B. Díaz et al. / Electrochimica Act

A

uncoat ed 5 nm 10 nm 20 nm 50 nm 5 nm 10 nm 20 nm 50 nm

10

10

10

Ccoat-fi tting Ccoat-theoreti cal ( Eq. 3)

Coa

ting

capa

cita

nce

/ F

Al2O3

250º C 160º C

B

uncoa ted 5 nm 10 nm 50 nm

10

10

10

Ccoat-fit ting Ccoat-theoretical ( Eq. 2)

Coa

ting

capa

cita

nce

/ F

Ta2O5

250º C

Fig. 7. Comparison of the experimental dielectric capacitance (obtained from theimpedance simulation shown in Figs. 4 and 5 and removing the contribution of then(

pbnfec

dirdtddttiMtTd

s

ative oxide capacitance) with theoretical capacitance (calculated from Eq. (2)) forA) the Al2O3 coatings and (B) the Ta2O5 coatings.

orosity values obtained from the LSV data) and the values obtainedy fitting (Ccoat, where the contribution of the native oxide waseglected) justifies the selection of the equivalent circuit employed

or the simulation. Fig. 7A shows the good agreement between thexpected capacitances and those obtained from fitting for the Al2O3oatings.

Although no quantitative analysis of the coatings in terms ofefects can be performed from the impedance data, some qual-

tative information can be extracted from the comparison of theesistance values. A higher resistance points to the growth of a lessefective coating. Thus, considering the coating resistance Rc ashe electric resistance through the existing pores and pinholes (asescribed in Fig. 6B), an increase of Rc, after taking into account theifferent coating thickness, is associated to a reduction in eitherhe size and/or density of these defects. Both at 160 and at 250 ◦C,he higher resistances obtained for the thicker Al2O3 coatings aren agreement with the lower porosities obtained from the LSV data.

oreover a better sealing (less defective) coating is confirmed byhe higher resistance obtained after growth at higher temperature.

he growth of a denser and more compact coating by ALD at highereposition temperature agrees with previous data [9,39].

The Bode plots corresponding to the Ta2O5 coatings are pre-ented in Fig. 5C. The impedance results show a worse performance

a 56 (2011) 10516– 10523

than for the Al2O3 coatings having the same thickness and thesame deposition temperature. For the Ta2O5 layers, only the thickercoating (50 nm) shows a marked improvement of the corrosionresistance. A phase angle lower than 90◦, indicating a less capacitivebehavior and thus a more corroding system [4,19], is observed forthe 20 nm coating. This shows an even higher degradation for thecoated sample than for the bare substrate after the immersion in theaggressive solution. Actually the equivalent circuit proposed abovehas no physical meaning and it cannot be employed in this particu-lar case. For the other samples, the fitting parameters are presentedin Table 2. Fig. 7B shows the comparison of the expected theoreticalcapacitance values (calculated from Eq. (2)) and the values obtainedby fitting (after correction following the Brug equation, Eq. (4), andremoving the contribution of the native oxide capacitance) for theTa2O5 coatings. A significant data dispersion can be observed com-pared to the Al2O3 coatings. A good agreement is obtained only forthe thicker Ta2O5 layer (50 nm). The poor and inconsistent resultsobtained with the thinner Ta2O5 layers imply some problems innucleation of Ta2O5 and thereby point to a need to improve the sur-face preparation. Whatever the reason for the nucleation problems,the Al2O3 process remains clearly much less affected. Indeed, theAl(CH3)3–H2O process is known as the best nucleating ALD processand has been successfully exploited in growing high quality Al2O3films also on polymer substrates [40].

The lowest values of the CPE power, n, obtained from the datafit for the Ta2O5 films are consistent with the large scatter betweenthe expected and the fitted capacitance values. In particular for the10 nm Ta2O5 coating, for which the lowest corrosion resistance wasobtained, the n parameter shows the highest deviation from theideal (pure capacitive) behavior (n < 0.9) among all samples, whichcorresponds to a higher heterogeneity of the coating. Comparingall samples, the highest n values, close to 1, corresponding to thecoatings having the lowest porosity values, were obtained with theAl2O3 films deposited at 250 ◦C. The better electrochemical perfor-mance obtained with these samples indicates less defects and/orpinholes relevant for corrosion. Moreover the parameter n close tounity shows that the surface is smoother and freer of defects, whichcorresponds to an ideal, higher corrosion resistant system. Liu et al.[19] have reported that a decrease of n value with an increase ofimmersion time of the coated sample was related to a lower capac-itive behavior since the coating becomes more conductive due tothe presence of a higher number of holes. Thus, and as confirmed byour data, the n parameter can be used as an indicator of the coatingquality and its performance.

4. Conclusions

Surface analysis by ToF-SIMS and electrochemical analysis in a0.8 M NaCl solution by LSV and EIS of ultra thin (5–50 nm) oxide(Al2O3 and Ta2O5) coatings deposited by ALD at 250 and 160 ◦C onstainless steel (316L) have been performed.

ToF-SIMS depth profiling allows the elemental characterizationof the bulk coating, interfacial and substrate regions. The bulk coat-ings are contaminated by traces of C and OH species originatingfrom unreacted fragments of the precursors used for deposition andtrapped in the films. A higher level of C contamination is observedin the Ta2O5 than in Al2O3 owing to more complete reactionsin the Al2O3 deposition. Increasing the deposition temperatureallows to decrease the C and OH contamination of the bulk coat-ing. The increase of C and OH contamination at the approach ofthe interfacial region points to an influence of the substrate surface

contamination in the first deposited layers. The interfacial regioncontains also a native oxide grown on the stainless steel surface,possibly modified by thermal oxidation prior to complete sealingby the coating.

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[[38] V.M.-W. Huang, V. Vivier, M.E. Orazem, N. Pébère, B. Tribollet, J. Electrochem.

B. Díaz et al. / Electrochimi

The potentiodynamic polarization data show that the depositedoatings can block pitting of the 316L substrate in the tested rangef potential (up to 0.5 V/SCE) in 0.8 M NaCl at room temperature.his is achieved with the 50 nm thick films deposited at 250 and60 ◦C. A drastic decrease of the current density is measured both

n the cathodic and in the anodic branches. The current reductionecomes more pronounced with increasing coating thickness andeaches four orders of magnitude in the best case (50 nm Al2O3eposited at 250 ◦C). The sealing performance of the coating wasssessed from the drop of the passive current and quantified ashe coating porosity, equivalent to the uncoated surface fraction.he porosity decreases with the thickness for all coatings, a betterealing performance being obtained with Al2O3 than with Ta2O5nd at a deposition temperature of 250 ◦C. The lowest porosity valuef 0.03% was achieved in this study for a 50 nm Al2O3 film.

The EIS Bode plots show one time constant for the coated andncoated systems. For the Al2O3 coatings, the experimental valuesf the coating capacitance obtained from the fit of the EIS data show

good agreement with values calculated assuming a flat parallellate capacitor. The n parameter of the constant phase elementsed for the evaluation of the capacitive behavior of the coatingfficiently allows for a qualitative ranking of the coatings in termsf presence of pores and defects, in line with the porosity data.he ranking of the coating resistances is also in agreement withhe classification obtained from the potentiodynamic polarization

easurements.The correlation between the porosity values extracted from the

V data and their ranking confirmed by EIS on one hand and theevel of C and OH traces observed by ToF-SIMS on the other handoints to a marked effect of the coating contaminants on the seal-

ng performance and on the corrosion resistance of the coatedystems.

cknowledgements

The research leading to these results has received fundingrom the European Community’s Seventh Framework ProgrammeFP7/2007-2013) under grant agreement n◦ CP-FP 213996-1 (COR-AL). Region Ile-de-France is acknowledged for partial support forhe ToF–SIMS equipment.

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