R1 revised version published in Applied Surface Science

The role of surface preparation in corrosion protection of copper with

nanometer-thick ALD alumina coatings

Shadi Mirhashemihaghighia, Jolanta Światowskaa, Vincent Mauricea,*, Antoine Seyeuxa, Lorena H. Kleina, Emma Salmib, Mikko Ritalab, Philippe Marcusa

aPSL Research University, CNRS – Chimie ParisTech, Institut de Recherche de Chimie Paris (IRCP), 11 rue Pierre et Marie Curie, 75005 Paris, France

bLaboratory of Inorganic Chemistry, Department of Chemistry, University of Helsinki, P.O. Box 55, FIN-00014 Helsinki, Finland

Abstract Surface smoothening by substrate annealing was studied as a pre-treatment for improving the

corrosion protection provided to copper by 10, 20 and 50 nm thick alumina coatings deposited

by atomic layer deposition. The interplay between substrate surface state and deposited film

thickness for controlling the corrosion protection provided by ultrathin barrier films is

demonstrated. Pre-annealing at 750°C heals out the dispersed surface heterogeneities left by

electropolishing and reduces the surface roughness to less than 2 nm independently of the

deposited film thickness. For 10 nm coatings, substrate surface smoothening promotes the

corrosion resistance. However, for 20 and 50 nm coatings, it is detrimental to the corrosion

protection due to local detachment of the deposited films. The weaker adherence of the

thicker coatings is assigned to the stresses accumulated in the films with increasing deposited

thickness. Healing out the local heterogeneities on the substrate surface diminishes the

interfacial strength that is bearing the stresses of the deposited films, thereby increasing

adhesion failure for the thicker films. Pitting corrosion occurs at the local sites of adhesion

failure. Intergranular corrosion occurs at the initially well coated substrate grain boundaries

because of the growth of a more defective and permeable coating at grain boundaries.

Keywords: Corrosion; Ultrathin coatings; Copper; Alumina; Atomic layer deposition; Surface treatment

* Corresponding author : Vincent Maurice (e-mail : vincent.maurice@chimie-paristech.fr)

1. Introduction

Along with the development of nanotechnology and device miniaturization, corrosion

protection of high precision components becomes increasingly challenging. For corrosion

protection by the use of coatings, methods for depositing ultrathin films have been studied

and continue to be developed. The major challenge for manufacturing ultrathin coatings for

corrosion protection is to overcome problems such as the presence of defects extending

through the whole film, poor conformality or continuity on morphological surface

heterogeneities and on chemical substrate heterogeneities, cracking of the coating and

impurities in the bulk coating or at the interface with the substrate. Significantly better

corrosion resistance being obtained with layers of micrometer thickness, another challenge is

to obtain similar protection with coatings of nanometer thickness which would reduce

dimensional modifications, costs and weight. To achieve high protection efficiency while

keeping the coating thickness small at the same time, the improvement of the existing

methods and the development of modern techniques for coating deposition are crucial.

Atomic Layer Deposition (ALD) is the most powerful technique for deposition of ultrathin

films that are conformal, continuous and pinhole-free even on high aspect ratio substrate

surfaces [1–4]. ALD has proven recently to be the main contender for deposition of

nanometer-thick oxide coatings for corrosion protection of metallic materials [5–21].

Even if ALD enables to precisely mimic the geometry of the substrate with high conformality,

thickness and composition control of the deposited films, the performance of ultrathin ALD

coatings in corrosion protection still depends on the surface state of the substrate

[17,18,22,23]. Easily detaching microparticles (e.g. remaining from surface preparation),

chemical substrate heterogeneities (e.g. second phase particles in alloys) and morphological

surface heterogeneities (including asperities or depressions of high aspect ratio) on the

substrate remain as possible sites of corrosion since, in case of exposure to the aggressive

2

environment via channel defects through the coating, they may trigger and/or accelerate

locally the corrosion process.

On copper [16,19,24] and other substrates [6,7,12,20], the sealing performance of nanometer-

thick alumina coatings prepared by ALD was demonstrated to be excellent. The outcome

picture is that the ALD alumina forms dense and amorphous barrier layers that are free of

crystalline defects but in which channel defects connecting the substrate to the aggressive

environment are present and responsible for residual corrosion that happens locally where the

substrate is defectively protected and exposed [6–8,17]. On copper exposed to corroding

aqueous electrolytes, corrosion and anodic dissolution currents could be decreased by up to

four orders of magnitude with 50 nm thick Al2O3 layers deposited on surfaces electropolished

in H3PO4 as a final surface treatment before the coating deposition [19]. Although remarkable

corrosion decrease was achieved, Atomic Force Microscopy (AFM) still evidenced a rather

rough copper surface under the coating as obtained by electropolishing. A high roughness of

the copper substrate has been proposed to be the reason for poorer electrical insulating

properties of the ALD Al2O3 films grown on this substrate [25].

Here we report on the tentative improvement of the corrosion protection of ALD alumina-

coated copper by surface smoothening and minimization of topographical heterogeneities as

obtained by substrate annealing prior to coating deposition. ALD alumina films were grown

to 10, 20 and 50 nm in thickness on substrates pre-annealed in reducing atmosphere. The

corrosion protection and its improvement were investigated in 0.5 M NaCl aqueous solution

by combining electrochemical analysis with Linear Sweep Voltammetry (LSV) and

Electrochemical Impedance Spectroscopy (EIS) and surface analysis by Atomic Force

Microscopy (AFM), Scanning Electron Microscopy (SEM) and Time-of-Flight Secondary Ion

Mass Spectrometry (ToF-SIMS).

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2. Experimental

Copper disks of 10 mm diameter, 2 mm thickness and 99.99% purity (from Goodfellow)were

used as substrate samples. The surface preparation started with mechanical polishing using

first 1200 abrasive SiC paper, and then water based diamond polishing suspensions of 6, 3, 1

and 0.25 μm, successively. After mechanical polishing, the samples were rinsed in successive

baths of acetone, ethanol and ultra-pure Millipore® water (resistivity > 18 MΩ cm).

Electropolishing followed immediately, performed in H3PO4 (60 wt.%) for 5 minutes at 1.4 V

versus a copper counter electrode. After electropolishing, the samples were rinsed with H3PO4

(10 wt.%) and ultrapure water successively, and then blow-dried with nitrogen. These

substrates were then annealed at 725°C for several hours under a flow of ultrapure

(99.9999%) hydrogen at atmospheric pressure, a procedure previously adopted to produce

well-crystallized single crystal surfaces [26].

The ALD layers were prepared in the Laboratory of Inorganic Chemistry of the University of

Helsinki. Before the coating deposition, the substrates were etched with 10 wt.% H3PO4 for 5

seconds, and then rinsed with ethanol, dried with compressed air and introduced in the reactor

in less than 5 minutes. This pre-treatment, adapted from Shimizu et al. [27], was applied in

order to remove the aged native copper oxide layer formed on the sample surface between the

surface preparation in Paris and the coating deposition in Helsinki. Deposition was done at

250°C using a Picosun SUNALE R-150 ALD reactor. Trimethylaluminium (TMA)

manufactured by Chemtura (AXION® PA 1300, purity 99.9%) and ultra-pure water

(resistivity > 18 MΩ cm) were used as precursors. More details on the deposition parameters

can be found in our previous study [19]. 120, 240 and 600 deposition cycles led to coating

thicknesses of 9.8, 20.3 and 51.4 nm, respectively. The coating thickness was measured with

X-ray reflectance (XRR Bruker AXS D8 Advance), employing a silicon wafer coated

simultaneously with the Cu substrates. The XRR curves were modeled with Leptos 7.05.

4

AFM imaging was performed in intermittent contact (tapping®) mode in air, using an Agilent

5500 microscope. A silicon cantilever with a resonance frequency in the range 200-400 kHz

and a force constant in the range 25-75 Nm-1 was employed. The silicon tip had a nominal

radius < 10 nm. SEM imaging was performed at LISE (Laboratoire Interfaces et Systèmes

Electrochimiques, CNRS - Université Pierre et Marie Curie) with a digital scanning electron

microscope S440 LEICA equipped with a tungsten field emission filament (FEG-SEM).

Secondary electron images were recorded at an accelerating voltage of 15 kV, a nominal

beam size of 1 nm and a measured beam current of 180 pA. Chemical imaging was performed

with a ToF-SIMS5 spectrometer (ION-TOF GmbH) operating at a residual pressure of 10-9

mbar. A pulsed 25 keV Bi+ primary ion source was employed, delivering 0.1 pA of current

over a 100 × 100 μm2 area analyzed with a resolution of 150 nm.

Before the electrochemical tests, the coated samples were rinsed in an ultrasonic bath of

ethanol for 5 minutes and dried with filtered compressed air. The bare sample was pre-treated

in 10 wt.% H3PO4 like prior to coating. A three-electrode cell was used as previously

described [19]. The electrochemical measurements were carried out at room temperature in a

0.5 M NaCl aqueous solution, prepared with ultra-pure Millipore® water and reagent grade

chemicals (NaCl Analar Normapur analytical reagent, VWR® BDH Prolabo®). The electrolyte

was bubbled with argon for 30 minutes prior to the measurements and during the analysis,

however without reaching complete deaeration as discussed previously [21]. The EIS

measurements were performed at OCP with an excitation signal set at 10 mV in a frequency

range of 10-2 to105 Hz. Fitting was done with Simad® software developed at LISE. The LSV

tests were performed after the EIS measurements with a scan rate of 1 mV s−1 from -0.8 V to

0.4 V vs SCE. An Autolab PGSTAT 30 potentiostat/galvanostat was used for cell control. The

working electrode area was 0.29 cm2 for all the electrochemical measurements.

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3. Results and discussion

3.1. Surface smoothening

Figure 1 presents the AFM images for the bare, 10, 20 and 50 nm ALD alumina-coated

samples prepared on the annealed copper substrates and comparative images for

electropolished but not annealed substrates coated likewise. The RMS roughness values are

presented in Figure 2 for both sample sets. On the electropolished substrates the surface

roughness is much higher and decreases with Al2O3 thickness. The smoothening with

increasing film thickness is due to the filling of holes on the substrate surface by the

conformal ALD Al2O3 film. On the annealed substrate smoothening is not detected because

these samples are smoother to start with.

Given that the same coatings were grown on both types of substrates, it can be concluded that

the surface morphology, quite different for the two sets of samples at this magnification, is

that of the substrates. The coating morphology is homogeneous in the presented images but

can be resolved at higher magnifications as previously shown [19]. The smoothening effect

brought by the annealing treatment is obvious. Roughness is mostly associated with the

presence of depressions (i.e. holes) created by the non-homogeneous dissolution process

during electropolishing. On the non-annealed substrate, the depressions have a large size

distribution and occupy a non-negligible surface fraction. On the annealed substrate, they

become smaller and cover a much lower surface fraction. Their depth (4 nm at the most) is

smaller than the deposited coating thickness, confirming that they correspond to substrate

depressions coated by ALD alumina. As expected, the annealing treatment cures the substrate

depressions by surface diffusion, and thereby smoothens the surface.

3.2. Corrosion protection

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The LSV polarization curves and EIS Bode plots obtained at OCP for the bare and coated pre-

annealed samples are presented in Figure 3.A and B, respectively.The polarization curves

have similar shapes to those previously obtained on the electropolished but not pre-annealed

samples and discussed together with the copper corrosion mechanisms in our previous paper

[19]. ALD alumina, being a dielectric insulator, has no electrochemical activity. As a result,

the polarization curves reflect the activity of the copper substrate and present a similar shape

for the coated specimens and a bare specimen, but with lower currents associated to the

residual electrochemical activity of the copper substrate exposed to the electrolyte by the

channel defects through the coatings. The parameters extracted from the polarization curves

are compiled in Table 1. It can be seen that the polarization resistance is higher for the coated

samples in comparison with the bare one, but that it decreases with increasing coating

thickness. This decrease was not expected since, in general, thicker coatings form better

barriers than thinner ones, as there is less possibility of electron transport and consequently

lower leakage currents through a thicker coating [25]. Furthermore, the densification of ALD

alumina with increasing deposited thickness has been previously concluded both from ToF-

SIMS and electrochemical measurements [6,7], also on the non-annealed copper substrate

[19].

Figure 4 shows the comparison of the polarization resistance and anodic current values for the

samples discussed here with those obtained on non-annealed substrates coated likewise [19].

For the 10 nm sample, the polarization resistanceis higher and anodic current lower on the

annealed substrate than on its non-annealed counterpart, indicating a beneficial effect of

surface smoothening on the corrosion resistance. Calculations of the uncoated surface fraction

or so-called coating porosity, based on Rp and ia values and as defined previously [19,28–31],

yield very close values for the annealed (PRp=0.83%, Pia

=1.26%) and non-annealed (

PRp=0.81%, Pia

=1.22 %) samples, showing that the coating sealing performance is

7

unchanged by the annealing pre-treatment. For the thicker 20 nm and 50 nm coatings, the

trend is reversed with lower polarization resistance and higher anodic current on annealed

substrates than on non-annealed counterparts (Figure 4). Besides, the corrosion protection

performance decreases with increasing coating thickness for the annealed substrates instead of

the expected increase indeed observed on the non-annealed substrates.The reason for this

unexpected performance on the annealed specimens is local detachment and loss of the

protection barrieras discussed in the following part of this paper.

Table 1 also shows that with increasing coating thickness the corrosion potential shifts

anodically and the cathodic Tafel slopes become more negative. The anodic shift of E corr is

consistent with a larger surface fraction exposing copper oxide to the electrolyte with

increasing coating thickness, which supports coating detachment for these samples. Also, the

larger│bc│ value implies higher cathodic activity near OCP. This is also consistent with a

larger surface fraction exposing copper oxide to the electrolyte. Indeed, it was discussed in

our previous work [21] and based on other studies [32–34] how the presence of Cu(I) species

on the copper surface promotes the oxygen reduction reaction dominating the cathodic

activity near OCP. In the conditions of our LSV tests, in which the potential is switched from

OCP to -0.8 V/SCE immediately before starting the anodic sweep, the native copper oxide,

present at the copper/ALD alumina interface exposed by the channel defects of the coating,

would only be partially reduced and the active surface could consist of Cu(I)/Cu(0) reported

to catalyse the reduction of residual dissolved oxygen present in our incompletely deaerated

electrolyte [21]. An increase of the surface fraction covered by such sites would increase the

cathodic activity, and thus the │bc│ value as measured here with increasing coating thickness.

On the EIS Bode plots (Figure 3.B), no separate time constants could be distinguished for the

coating and the susbtrate, which is assigned to their overlap. A hint of an extra time constant

can be distinguished at the lowest frequencies in the phase angle plots of the 20 and 50 nm

8

samples. These points correspond to anodic mass transport limitation via the copper oxide

layer covering the substrate [35]. The anodic mass transport has been attributed to the

diffusion of CuCl2- away from the electrode, through both the Cu2O layer and the electrolyte

[36,37]. Observation of the anodic mass transport in the Bode plots for these samples and not

for the 10 nm sample is again in agreement with the detachment of the thicker coatings and

consequent exposure of a larger surface fraction of the copper oxide-covered substrate to the

electrolyte.

The impedance results were fitted with the equivalent circuit presented inFigure 3.C [19,35].

Rcta and Wc, representing the parallel combination of anodic and cathodic faradic responses at

OCP, are the anodic charge transfer resistance and the cathodic Warburg impedance,

respectively. Wc represents the oxygen diffusion limitation. CPE is the constant phase element

representing the capacitive response of a double layer or an oxide. Details about the choice of

the equivalent circuit were discussed previously [19,35]. The anodic mass transport related to

the few points at low frequencies was not taken into account in the equivalent circuit due to

the low resolution of this loop in agreement with what was done by Torres et al. [35]. The

same equivalent circuit was used for the bare and coated samples, as no distinguishable time

constants could be recognized for the coating and substrate in the Bode plots (Figure 3.B).

The fitting results are compiled in Table 2.

The anodic charge transfer resistance values obtained from the fitting match well with the

diameters of the depressed semicircles extrapolated to the real axis (Zj = 0) in the Nyquist

graphs (not shown here). The Rcta values are higher for the coated samples in comparison with

the bare one, but again they decrease with increasing coating thickness in agreement with

what was observed from the polarization resistances (LSV results).

The Warburg coefficient (S) is higher for the coated substrates in comparison with the bare

one in agreement with increased length of oxygen diffusion for the coated specimens as

9

discussed previously [20,21]. However, S decreases with increasing coating thickness, which

is again a result of the increase of the surface area exposed to the electrolyte due to the

coating detachment and consequent lower oxygen diffusion limitation.

3.3. Failure analysis of the thicker coatings

SEM images of the 50 nm ALD alumina-coated pristine sample are displayed in Figure 5. At

large scale, the surface appears polycrystalline with grains delimited by grain boundaries and

triple joints (Figure 5.A,B). This morphology arises from the crystalline structure of the

copper substrate under the conformal and amorphous alumina coating. On the grains, the

coating presents two types of defects. One type has a circular shape with the coating lifted-up

but still present as magnified in Figure 5.E, suggesting local deadhesion of the coating. At low

magnification (Figure 5.A,B), many of these defects can be found with varying sizes,

suggesting different stages of deadhesion propagation. In the other type of defects, the coating

is missing. Detached pieces of the coating can be seen as flakes in or at the periphery of the

region where the coating is lost (Figure 5.B,D). Flakes can also be seen away from a missing

coating defect (Figure 5.C). At low magnification (Figure 5.A,B), the lost coating defects are

fewer and larger than the lifted-up coating defects, suggesting that the loss of the coating

could be a stage subsequent to deadhesion. The grain boundaries and triple joints appear well

coated (Figure 5.B,C), as expected from the excellent conformality of ALD films.

Figure 6 shows ToF-SIMS chemical maps of the same 50 nm ALD alumina-coated pristine

sample. The analyzed region contains two grain boundaries, one triple joint and one lost

coating defect. It confirms the SEM observations. The lost coating defect has a nearly circular

shape with a diameter of about 20 µm, consistent with the shape and dimensions of the

smallest lost coating defects observed by SEM (Figure 5.A). The loss of the coating is proven

by the map of the AlO2- ions that are characteristic of the coating and have no intensity inside

10

the defect. In the O- map, a dark ring is observed instead of a dark disk. The intensity inside

the ring is a result of oxygen coming from copper oxide covering the substrate surface

exposed by the loss of the coating. The presence of copper oxide is confirmed by the CuO -

and Cu- maps. The Cl- and C- maps confirm the chlorine and carbonaceous contamination

usually observed on ALD alumina-coated surfaces as discussed previously [6,7,9,15,24,38].

No marked variation of this contamination is observed in the coating defect exposing the

substrate. At the grain boundaries and triple joint, there is no evidence of missing or defective

coating at the resolution of the ToF-SIMS images.

Figure 7 shows SEM images of the 20 nm ALD alumina-coated sample taken after the anodic

polarization test. As seen in Figure 7A, pits have formed both on the grains and at grain

boundaries. They result from localized corrosion occurring in the defects of the ALD coating.

The rest of the copper surface remained non-attacked and well-protected by the non-defective

coating. Interestingly, some grains in the upper part of the imaged region better resist to

localized corrosion by pitting as a result of the formation of a less defective and better

adherent coating. In the corroded grains, the less numerous and larger pits most likely

correspond to the defects where the coating was initially lost. The more numerous and smaller

pits presumably formed in the defects where the coating was initially lifted-up as a result of

the local deadhesion. The formation of pits in this type of defects suggests breakdown and

partial and/or complete loss of the lifted-up coating upon immersion at OCP as supported by

the EIS data.

Intergranular corrosion is also observed at grain boundaries (Figure 7.B,C), meaning that

corrosion has propagated in depth at the grain boundaries although they seemed initially well

coated. The nearby grain surface is not attacked. It is proposed that, upon immersion, the

electrolyte has penetrated the coating and reached the copper substrate via channel defects of

the more permeable coating formed at the grain boundaries of the substrate. Then, upon

11

anodic polarization, dissolution has occurred faster at grain boundaries because of preferential

intergranular corrosion. As a result, the coating-substrate interface was trenched along the

grain boundaries with liberation of corrosion products which caused the coating to swell

(Figure 7.B) and to fracture in some parts leaving remnants of the deposited coating (Figure

7.C). Thus, although the grain boundaries seemed to be initially well coated (Figure 5.C), the

substrate sealing in these sites appears more defective and insufficient to provide efficient

protection in the conditions of severe dissolution imposed by anodic polarization.

Thus, it appears that relatively weak adherence of the coating is the main cause for the locally

poor corrosion protection of the substrate, leading to coating breakdown and localized

corrosion by pitting in aggressive conditions. Weak adherence results in the local detachment

of the coating and is observed at grains already prior to immersion. The EIS data confirm that

the weak adherence causes the corrosion protection failure after immersion at OCP, except for

the more strongly adherent 10 nm coating. Preferential intergranular corrosion of the coated

specimen is another type of degradation observed in conditions of severe anodic dissolution.

It does not result from an initial poorer adhesion of the coating at grain boundaries but rather

from the growth of a more defective and permeable coating at grain boundaries.

The reason for the weaker adherence of the ALD alumina coatings on the annealed copper

substrate as compared to the non-annealed substrate could be the excessive smoothness of the

substrate surface. Although reducing the roughness is important for the integrity and fatigue

performance of ultrathin coatings, sometimes a minimum surface roughness may be necessary

for good coating adhesion [39,40]. An initial roughness has been shown to increase the

interface area to interlock the coating with the substrate [39]. If such a mechanical effect is

valid in the present system, it should benefit from stronger interlocking provided by the more

numerous and larger depressions found on the electropolished copper surface but healed out

from the annealed substrate by surface diffusion. Surface segregation of impurities could be a

12

factor contributing to a weaker chemical bonding, and thus to the coating detachment at the

interface between coating and pre-annealed substrates. However, this is not supported by our

electrochemical data that evidence the same sealing performance of the 10 nm coating

independently of the substrate pre-annealing, and thus the absence of detachment of this

thinner coating deposited on the same pre-annealed substrate as the thicker detaching

coatings. Besides, an excess of segregated impurities (i.e. sulfur) is not supported by ToF-

SIMS data that show the same depth profiles on pre-annealed substrates and their

electropolished counterparts [19].

The decrease of the adhesion of ALD alumina coatings with increasing coating thickness has

been observed with adhesion tests by Marin et al. [41]. Furthermore, in the work of Miller et

al. [42], it has been shown that the mechanical robustness of ALD alumina coating increases

with decreasing coating thickness from 125 nm to 5 nm. The change of the mechanical

properties is the result of the increased stresses accumulated in the deposited oxide films with

increasing thickness [43]. In the present work, the weaker adherence of the thicker coatings

may also result from the stresses accumulated in the films. It is suggested that healing out the

local heterogeneities on the substrate surface by annealing diminishes the interfacial strength

that is bearing the stresses of the deposited films, thereby increasing adhesion failure for the

thicker films.

Apart from the mechanical properties of the ALD coating/substrate interface,

electrochemically induced delamination may assist detachment and breakdown of the thicker

coatings. Indeed, the cathodic oxygen reduction reaction occurring at the coating-substrate

interface exposed by the channel defects leads to a local increase of pH triggering local

dissolution of alumina at the interface with the substrate, and thus interfacial trenching and

local delamination of the coating [44]. It is suggested that the longer diffusion length through

the thicker coating could retard the diffusion of OH- ions away from the interface to the bulk

13

electrolyte and thus result in a faster local increase of pH at the coating-substrate interface,

and therefore faster trenching and local delamination of the alumina coating. Thus, even if the

primary reason for the local coating detachment appears to be a consequent effect of the

surface smoothness, increasing with the deposited thickness, the coating detachment,

breakdown and loss are assisted upon immersion in aqueous solution by an electrochemically-

induced delamination mechanism, including oxygen reduction at or near OCP and anodic

dissolution of corrosion products at anodic potentials.

4. Conclusions

Electrochemical and surface analysis were applied to study the effect of surface smoothening

by substrate pre-annealing on the corrosion protection provided to copper by 10, 20 and

50 nm thick ALD alumina coatings. The conclusions are the following:

Substrate pre-annealing at 725°C was effective to reduce the surface roughness of the

coated specimen to less than 2 nm (RMS value) independently of the deposited film

thickness, healing out the dispersed topographic heterogeneities left by

electropolishing.

For 10 nm coatings, the corrosion protection is excellent with polarization and charge

transfer resistances and active dissolution current increasing and decreasing by two

orders of magnitude, respectively, as compared with the bare substrate. Substrate

annealing slightly improves the corrosion resistance by reducing the local topographic

heterogeneities, however without affecting the sealing performance of the 10 nm ALD

barrier film.

For 20 and 50 nm coatings, substrate annealing is detrimental to the corrosion

protection provided by the coatings. With increasing coating thickness, the

polarization and charge transfer resistances decrease and the active dissolution current

14

increases instead of increasing and decreasing, respectively. Local coating detachment

from the annealed substrate, inferred from both LSV and EIS analysis, was confirmed

by FEG-SEM and ToF-SIMS imaging before immersion in the electrolyte.

The weaker adherence of the thicker coatings is assigned to the stresses that

accumulate in the films with increasing thickness. It is suggested that healing out the

substrate surface local heterogeneities by annealing diminishes the interfacial strength

that is bearing the stresses of the deposited films, thereby increasing adhesion failure

for the thicker films.

Pitting corrosion occurred at the local sites of adhesion failure. A local increase of pH

at or near OCP as a result of cathodic oxygen reduction is suggested to promote the

detachment of the thicker coatings more permeable in these sites. Intergranular

corrosion occurred at the initially well coated substrate grain boundaries, also

suggesting the local formation of more permeable coatings at these substrate sites.

These results provide further evidence of the key role played by the substrate surface

preparation for controlling the corrosion protection provided by ultrathin deposited films. It is

shown that the interplay between surface state and film thickness must be considered for

optimizing the corrosion resistance.

AcknowledgementsRegion Ile-de-France is acknowledged for partial funding of the ToF-SIMS equipment. This

work is linked to the Finnish Centre of Excellence in Atomic Layer Deposition funded by the

Academy of Finland.

15

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Table 1. Parameters related to LSV polarization curves in Figure 3.A for the bare and 10, 20

and 50 ALD alumina-coated annealed copper substrates

Ecorr / V (SCE) ia (at E = -0.2 V/SCE) / A.cm-2 RP / Ω.cm2 ba / V.dec-1 bc / V.dec-1

Bare substrate -0.359 1.19 × 10-5 2.61 × 105 0.070 -0.215

Al2O3 10nm -0.356 1.50 × 10-7 2.90 × 107 0.065 -0.350

Al2O3 20nm -0.336 1.52 × 10-6 1.68 × 106 0.064 -0.375

Al2O3 50nm -0.308 2.71 × 10-6 3.82 × 105 0.065 -

Table 2. Parameters for the EIS data in Figure 3.B fitted with the equivalent circuit shown in

Figure 3.C for the bare and 10, 20 and 50 ALD alumina-coated annealed copper substrates

Re / Ω.cm2 Rcta / Ω.cm2 Q / F.cm-2.s(α−1) n

S / Ω.s-0.5.cm2

(ZW= S

√ jω2

)

Bare substrate

11.28 ± 0.10 (1.01± 0.20) × 105 (3.70 ± 0.12) × 10-5 0.91 ± 0.00 (1.03 ± 0.01) × 104

Al2O3 10nm 11.22 ± 0.58 (6.11 ± 3.35) × 106 (7.80 ± 0.23) × 10-7 0.98 ± 0.01 (9.34 ± 3.06) × 106

Al2O3 20nm 11.70 ± 0.17 (5.31± 1.46) × 105 (1.28 ± 0.46) × 10-6 0.98 ± 0.01 (6.14 ± 3.61) × 105

Al2O3 50nm 11.88 ± 0.48 (1.81± 0.27) × 105 (2.60 ± 0.36) × 10-6 0.95 ± 0.00 (1.64 ± 0.35) × 105

18

1 µm 1 µm

1 µm

1 µm

1 µm 1 µm

10 nm 20 nm 50 nmA

nnea

led

subs

trat

eE

lect

ropo

lishe

d su

bstr

ate Δz = 17 nm Δz = 5 nm Δz = 6 nm

Δz = 66 nmΔz = 46 nmΔz = 70 nm

Bare

1 µm

1 µmΔz = 66 nm

Δz = 29 nm

Figure 1. Topographic AFM images (5 × 5 µm2) for bare samples and 10, 20 and 50 ALD

alumina films deposited on annealed Cu substrates (upper row) and electropolished Cu

substrates (lower row)

19

Electropolished Annealed

02468

101214161820

BareBare

Rou

ghne

ss (R

ms)

/ nm

50nm50nm 20nm20nm 10nm10nm

Figure 2. Roughness values (RMS) obtained by AFM for bare samples and 10, 20 and 50

ALD alumina films deposited on electropolished or annealed Cu substrates

20

-0.8 -0.6 -0.4 -0.2 0.0 0.2 0.410-12

10-11

10-10

10-9

10-8

10-7

10-6

10-5

10-4

10-3

10-2

10-1

Cur

rent

den

sity

/ A

cm

-2

Potential / V (SCE)

Bare Substrate 10nm Al2O3

20nm Al2O3

50nm Al2O3

10-2 10-1 100 101 102 103 104 105100

101

102

103

104

105

106

107

108

109

Bare substrate10nm Al2O3 20nm Al2O3 50nm Al2O3

Z

/ .c

m2

fHz

0

10

20

30

40

50

60

70

80

90

Phas

e an

gle

/ deg

ree

A

B

C

W c

CPE

Re Rct

Figure 3. (A) LSV polarization curves, scan rate of 1 mV.s−1 and (B) EIS Bode plots at OCP

for bare and 10, 20 and 50 ALD alumina-coated annealed copper substrates in 0.5 M NaCl

under argon bubbling (pH = 6). (C) Equivalent circuit used to fit the EIS data for the bare and

coated copper substrates

21

10-9

10-8

10-7

10-6

10-5

10-4

i a / A

.cm

2

Annealed substrate Electropolished substrate

10 nm 20 nm 50 nm

105

106

107

108

109

1010

Rp /

Ohm

.cm

2

Annealed substrate Electropolished substrate

10 nm 20 nm 50 nm

A

B

Figure 4. Comparison of polarization resistance (A) and anodic currents (at E = -0.2 V/SCE)

(B) of the ALD alumina coated samples prepared on annealed and electropolished copper

substrates. Data for electropolished copper are from Ref. [19]

22

100 µm 20 µm

2 µm1 µm 1 µm

A B

C D E

Figure 5. FEG-SEM images for the pristine 50 nm ALD alumina coated sample prepared on

the annealed copper substrate

23

AlO2-

63Cu- 63CuO- 63Cu2-

O-

C-

Cl-

Figure 6. ToF-SIMS chemical images (100 × 100 µm2) for the pristine 50 nm ALD alumina

coated sample prepared on the annealed copper substrate

24

100 µm

A

10 µm

B

2 µm

C

Figure 7. FEG-SEM images for the 20 nm ALD alumina coated sample prepared on the

annealed copper substrate and taken after anodic polarization in 0.5 M NaCl

25