Electronic properties of thermally formed thin iron oxide films

Electrochimica Acta 52 (2007) 7617–7625

Electronic properties of thermally formed thin iron oxide films J. Wielant a,∗,1 , V. Goossens a , R. Hausbrand b , H. Terryn a a

Department of Metallurgy, Electrochemistry and Materials Science, Vrije Universiteit Brussel, Pleinlaan 2, B-1050 Brussels, Belgium b Arcelor Research Industry Gent-OCAS nv, John Kennedylaan 3, B-9060 Zelzate, Belgium Received 17 July 2006; received in revised form 12 October 2006; accepted 17 December 2006 Available online 21 January 2007

Abstract The oxide layer, present between an organic coating and the substrate, guarantees adhesion of the coating and plays a determinating role in the delamination rate of the organic coating. The purpose of this study is to compare the resistive and semiconducting properties of thermal oxides formed on steel in two different atmospheres at 250 ◦ C: an oxygen rich atmosphere, air, and an oxygen deficient atmosphere, N2 . In N2 , a magnetite layer grows while in air a duplex oxide film forms composed by an inner magnetite layer and a thin outer hematite scale. The heat treatment for different amounts of time at high temperature was used as method to sample the thickness variation and change in electronic and semiconducting properties of the thermal oxide layers. Firstly, linear voltammetric measurements were performed to have a first insight in the electrochemical behavior of the thermal oxides in a borate buffer solution. Electrochemical impedance spectroscopy in the same buffer combined with the Mott–Schottky analysis were used to determine the semiconducting properties of the thermal oxides. By spectroscopic ellipsometry (SE) and atomic force microscopy (AFM), respectively, the thickness and roughness of the oxide layers were determined supporting the physical interpretation of the voltammetric and EIS data. These measurements clearly showed that oxide layers with different constitution, oxide resistance, flatband potential and doping concentration can be grown by changing the atmosphere. © 2007 Elsevier Ltd. All rights reserved. Keywords: Thermal oxide film; Impedance spectroscopy; Ellipsometry; Atomic force microscopy; Mott–Schottky concept

1. Introduction Since many years, steel and galvanised steel substrates have been covered by organic coatings to improve the corrosion protection efficiency of the easily corroding substrates. Different factors play an important role in the protection efficiency of these systems. Firstly, the chemical composition and constitution of the organic coating [1–3]. Secondly, the adhesion of the organic coating through interfacial bonds between the coating and the thin oxide layer on top of the metallic substrate [4–11]. And, finally, the morphological and electrical properties of the oxide layer [7,12–14]. In corrosive environments, coatings will often fail due to delamination. Because of the local electrochemical behavior of the delamination process, the delamination rate is related to the electron transfer reactions taking place between the anodic and cathodic sites. For steel, which is subject to cathodic

∗ 1

Corresponding author. Fax: +32 2 6293200. E-mail address: [email protected] (J. Wielant). ISE member.

0013-4686/$ – see front matter © 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2006.12.041

delamination [15–19], it is favourable to slow down the electron transfer rate between the anodic sites, e.g. scratches, and the cathodic sites, e.g. the delamination front of the organic coating. Since the electrons reaching the cathodic site travel through the oxide layer, the oxide resistance is one of the crucial parameters to scale down the delamination rate. From semiconductor physics and electrochemistry, it is known that near the flatband potential the oxide conductivity varies with orders of magnitude. So, the position of the flatband potential relative to the electrode potential gives valuable information about the electrochemical behavior of the interfacial oxide layer [20]. Both iron oxides and non-ferrous oxides are used as interfacial oxide films. Interfacial oxides can be native [21], synthesised [22,23] or deposited by, e.g. CVD [24]. Iron oxides can be produced in several ways, such as electron beam deposition [25], sputter deposition [26,27], alkaline treatments [28,29], electrochemically [29–31] and also thermally [32]. The thermal way of oxide formation involves an oxidation/reduction reaction occurring in the dry state. The reaction, limited by the availability of oxygen, takes place in the growing oxide layer [33]. The existence of many thermal iron oxides with

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different structures and properties make interpretation of these oxide layers very difficult [34]. When heated to a temperature of 200–300 ◦ C in atmospheric air, well-defined duplex films are formed with an inner core of magnetite (Fe3 O4 ) and an outer layer composed of hematite (␣-Fe2 O3 ) [32,35–37]. In an oxygen deficient atmosphere, like N2 , the formation of a pure magnetite scale is promoted [32]. Magnetite and hematite have completely different structures. The spinel structure of magnetite consists of oxide ions in cubic closed-packed arrays with cations in the octahedral and tetrahedral interstices [38,39]. The Fe3+ cations occupy octahedral and tetrahedral sites, while the remaining octahedral positions are occupied by Fe2+ cations. Hematite, on the other hand, exhibits a hexagonal close packing of oxygens with mainly Fe3+ cations distributed in the octahedral interstices [32]. Mostly, magnetite and hematite behave like n-type semiconductors due to point defects in the oxide structure, like Fe2+ ions in the Fe3+ matrix and O2− vacancies. It is well known that these defects play an important role in the corrosion of metals, influencing the transport through the oxide layer [40,41]. The conductivity of hematite is very low compared to magnetite because of the very low Fe2+ donor state density [32]. Senkevich et al. [42] characterized oxide films formed at 500 ◦ C by ‘dry’ impedance spectroscopy. The outer atomic layers of iron oxide films in contact mode, e.g. in contact with an electrolyte, will behave totally different from the inner bulk oxide. In these conditions, the electronic charge is redistributed at the semiconductor/solution interface resulting in the formation of a space charge double layer in the outer oxide layer and a Helmholtz double layer in the electrolyte near the oxide surface [20]. By varying the potential at the semiconductor electrode, a correlation can be found between the total capacitance of the semiconductor/electrolyte system and the applied potential on the working electrode. Since many years, this approach, the Mott–Schottky concept, has been demonstrated to be a useful tool for the characterization of semiconductor properties, in combination with electrochemical impedance spectroscopy (EIS) [43–45]. This approach has been used successfully to investigate passive oxide layers immersed in borate buffers [26,46,47]. The purpose of this study is to compare the resistive and semiconducting properties of thermal oxides formed on steel in two different atmospheres at 250 ◦ C: an oxygen rich atmosphere, air, and an oxygen deficient atmosphere, N2 . Firstly, linear voltammetric measurements were performed to have a first insight in the electrochemical behavior of the thermal oxides in the borate buffer solution. Next, electrochemical impedance spectroscopy was used to characterize the semiconducting properties of the formed oxide layers. By spectroscopic ellipsometry (SE) and atomic force microscopy (AFM), respectively, the thickness and roughness of the oxide layers was determined supporting the physical interpretation of the voltammetric and EIS data. In a next paper, the role of the interfacial oxide films and the influence of the electronic oxide properties on the corrosion resistance efficiency and delamination rate of organic coatings will be studied more in detail.

2. Experimental conditions 2.1. Sample preparation The used substrate is an interstitial free DC06 steel according to the EN10130(98) norm. The substrate was cut in a circular shape, with a diameter of 3 cm and a thickness of 1 mm. In order to reduce the surface roughness, the surface was mechanically polished with SiC-paper (500–800–1200–4000 grid) and finished with 1 ␮m grade diamond paste to obtain mirror-like surfaces. To form an iron oxide layer, samples were treated in an oven at 250 ◦ C under an oxygen rich atmosphere, air, or under an oxygen deficient atmosphere, N2 (AlphagazTM 2 grade (O2 < 0.1 ppm, H2 O < 0.5 ppm) supplied by Air Liquide). After the thermal treatment, the samples were cooled down to ambient temperature in open air and N2 , respectively. Finally, the samples were cleaned in acetone and chloroform for 5 min to remove the organic contamination as much as possible. 2.2. Spectroscopic ellipsometry The ellipsometric experiments were performed using a J.A. Woollam Co. VASE (Variable Angle Spectroscopic Ellipsometer) working in the UV–vis–NIR (0.98–4.98 eV) spectral range with an energy resolution of 0.1 eV at three angles of incidence, 60◦ –65◦ –70◦ . The ellipsometric data were analysed with the WVASE32 Version 3.441 software developed by J.A. Woollam Co. The optical constants and thickness of the thermally formed oxide layers were determined using the optical modelling procedure described below. A more detailed description of this procedure can be found in literature [48]. The optical constants of the substrate were obtained in situ after reduction of the steel sample at −1 V for 30 min in borate buffer solution and inserted directly into the model. In literature, no optical data for the magnetite film have been found. So, the data obtained from the oxide layers formed in N2 were used to model the magnetite layer with an oscillator model, composed of three Tauc–Lorentz oscillators and one Drude oscillator. For the hematite layer, the optical constants found in literature [49] were used as a reference to construct an oscillator model consisting of two Tauc–Lorentz oscillators. In this case, the data from the thermal oxide layers formed in air were fitted to an optical model containing both oxide layers. 2.3. Linear voltammetry (LV) and electrochemical impedance spectroscopy (EIS) For the voltammetric measurements and impedance measurements, a PGSTAT-30 from Autolab, Ecochemie B.V. was used. The electrochemical cell contained a silver/silver chloride reference electrode (Ag/AgCl), a large platinum grid as counter electrode and the oxide covered steel sample as working electrode. The exposed area was approximately 0.95 cm2 . The voltammetric scans were obtained at room temperature and open to air in a borate buffer solution (0.075 M Na2 B4 O7 ·10H2 O + 0.3 M H3 BO3 ; pH 8.2). The scanned potential ranged from −0.4 to 1.2 V versus the reference electrode at a

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scan rate of 1 mV s−1 . The electrochemical impedance measurements were performed in the 10 kHz–10 mHz frequency range with 10 frequency points per logarithmic decade. In order to reduce the phase shift induced by the reference electrode in the high frequency range (>105 Hz), a Pt wire coupled with a capacitance of 10 nF was placed in parallel with the reference electrode [50]. The amplitude of the sinusoidal voltage was 10 mV rms. The impedance spectra were obtained at ambient temperature and open to air in the borate buffer solution. The EIS measurement time was approximately 15 min. The quantitative values for the different components in the equivalent electric circuit were derived by the Complex Nonlinear Least Squares software developed by R. Macdonald. For the Mott–Schottky analysis, the measurements were performed at different dc potentials. The potential ranged from −0.4 to 1.0 V versus Ag/AgCl. The potential step size was 100 mV. The polarisation time before each measurement was 30 s. The voltammetric curves and impedance measurements were obtained after 1 h of immersion of the sample in the electrolyte. 2.4. Atomic force microscopy An atomic force microscope (Nanoscope IIIa; Digital Instruments; VEECO), equipped with a piezoscanner (E scanner) and silicon nitride tip, was used to image each sample at room temperature open to air. The images (512 × 512 points) were obtained in Tapping mode (TM). The projected area of the imaged surface was 1 ␮m × 1 ␮m. The RMS roughness and surface area of the studied samples were calculated by using the Nanoscope 6.11 software. 3. Results and discussion 3.1. Voltammetry studies Several steel samples were treated for different amounts of time in an oven in the presence of air or N2 . After oxide formation in both atmospheres the samples were characterized by linear voltammetry in a borate buffer. Prior to the voltammetric experiment, the open circuit potential was recorded after 1 h of immersion. The obtained open circuit potentials versus Ag/AgCl are shown in Fig. 1. The OCPs of the oxides formed in N2 fluctuate around −180 mV. The air formed oxides show OCP-values at approximately −80 mV for samples treated shorter than 8 min and around −160 mV for longer times. Fig. 2 presents the polarisation curves of several thermal oxides after 1 h of stabilisation along with the polarisation curve of the polished steel substrate (after reduction for 30 min in the solution). The difference in current densities measured in the cathodic potential region and at potentials above 0.8 V is very small for the thermal oxide layers. In the cathodic region, H2 evolution and oxygen, Fe3+ and Fe2+ reduction processes take place [51]. For the reduced polished steel sample, the cathodic current is lowest as a result of the absence of divalent and trivalent iron. Above 0.8 V, the anodic current starts to increase due to O2 evolution. On the other hand, large differences in the anodic potential region between OCP and

Fig. 1. Open circuit potential vs. Ag/AgCl of different thermal oxide layers as a function of heat treatment time in N2 (a) and air (b) at 250 ◦ C. The OCPs are obtained after 1 h of immersion in borate buffer (pH 8.2).

0.8 V are recorded. All samples behave passive in this potential region. The passive current is highest for the polished steel sample. From literature, it is known that in the used conditions a highly defective ␥-Fe2 O3 /Fe3 O4 layer is formed in the passive

Fig. 2. Voltammetric curves of polished steel (as reference) and thermal oxides formed after treating in N2 and air for different amounts of time. The measurements are performed in borate buffer (pH 8.2) after 1 h of immersion. The scan rate is 1 mV s−1 .

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Fig. 3. Variation of SE thickness of the oxide layer(s) formed in N2 (a) and in air (b) at 250 ◦ C as a function of heat treatment time.

region [31]. For all thermal oxides, the anodic current density is much lower than for the polished steel sample meaning strongly limited dissolution of the iron substrate occurs. Moreover, compared to the thermal oxides formed in N2 , the air formed oxide layers reduce even more the anodic current. The lowest anodic currents were measured on the samples treated for 5 and 8 min in air. The anodic part of the polarisation curve causes the OCP shift in anodic direction. In Fig. 3, the thickness of the oxide layers formed in air and N2 as a function of time in the oven is shown. In the presence of air, a thin hematite layer is present with varying thicknesses between 2.5 and 7 nm. For oven times shorter than 8 min, a slow growth of the hematite layer from 4 to 7 nm is noted. At 8 min, the hematite layer reaches its maximum thickness followed by a sharp thickness drop between 8 and 15 min. For longer amounts of time, the hematite thickness constantly remains around a value of 2.5 nm, while the magnetite layer is growing further. This may indicate that a part of the hematite layer is converted from hematite into magnetite, probably because of an oxygen deficiency in the layer. With increasing layer thickness, the diffusion rate of the oxygen decreases, thereby causing the iron to oxidise incompletely and forming magnetite instead of hematite. As can be seen, the magnetite layer grows linearly with the logarithm of time in both cases. In N2 , however, the magnetite layer grows at higher rate.

This can be explained by the oxidation/reduction reactions and electronic/ionic transport taking place during the oxide growth [32]. For the oxides grown in air, the present hematite layer at the oxide surface highly hinders the transport rate of electrons and ions through the oxide resulting in a slower oxide growth. Bringing Figs. 1–3 together, it can be concluded that the lowest passive current densities were measured on the samples with the thickest hematite layers and the OCP values are directly related to the hematite thickness. The magnetite film thickness does not affect directly the passive current flowing through the oxide layer. If it is assumed that in the bulk film high-field assisted migration of defects proceeds, then the anodic current is related to the transport rate of ionic species through the oxide layer [52]. Considering magnetite contains much more defects or dopants than hematite layers [32], the ionic transfer rate will be determined by the hematite layer and the potential drop over the magnetite layer can be neglected compared to the drop over the hematite film. Consequently, for equal anodic potentials, the field strength in the hematite film is inversed proportional to the hematite thickness. Thus, the oxide films with the thickest hematite layer show low passive current densities compared to oxides with a thinner hematite layer. The formation of space charges at the semiconductor/electrolyte interface and their influence on the distribution of the potential drop in the system is not taken into account in the present treatment of the ion transport. The presence of the different thermal oxide layers has a limited effect on the cathodic reactions and O2 evolution rate. This means that in these reactions no ionic transport through the oxide film is involved. It is clear that these reactions take place at the oxide/electrolyte interface instead of the metal/oxide interface like the iron dissolution reaction [51,53]. 3.2. Characterisation of the oxide films with EIS Electrochemical impedance spectroscopy was used to evaluate the semiconducting properties of the thermal oxide layers. Mott–Schottky plots were built from parameters obtained from impedance data modelling. In order to construct the Mott–Schottky curves, impedance measurements were performed at potentials between −0.4 and 1.0 V. The lower potential limit was chosen to minimise oxide reduction. In Fig. 4, the bode plots performed at some selected potentials of a polished steel sample treated for 8 min in air are shown. As can be seen from these spectra, the total impedance of the oxide/electrolyte system depends on the applied potential. A higher potential results in a higher impedance. The phase diagram of the Bode plot shows an asymmetric shape due to the presence of two double layers at the semiconductor/electrolyte interface: the space charge layer and the Helmholtz layer. No capacitance contribution by the bulk oxide has been taken into account because the appropriate description of the capacitance values is provided by a semiconductor and not by a dielectric approach [26]. The equivalent circuit used to fit the experimental data consists of two Voight elements in series with the electrolyte resistance Re . The circuit is presented in Fig. 5. The first Voight element significates the total double layer, which includes the

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capacitance could be calculated as follows [26,46]: Csc =

Fig. 4. Bode plot of thermal oxide layer formed after 8 min in air. Spectra are obtained at different potentials vs. Ag/AgCl in borate buffer (pH 8.2) after 1 h of immersion.

−1 (f/f0 2π)nsc Z sinφ

(1)

where nsc is the exponent of the CPE attributed to the space charge layer, f the evaluation frequency, Z the EIS impedance value, φ the phase at frequency f and f0 is the reference frequency, arbitrarily set to 1 Hz. For the studied oxides, the Helmholtz double layer always contributed to the measured impedance value, even at high frequencies. In order to obtain reliable values for Csc , these contributions, known from the data fitting, had to be subtracted from the obtained EIS data. By doing so, for a steel sample treated for 8 min in air, the Csc value at 0.2 V was 1.24 × 10−5 and 1.22 × 10−5 F at, respectively, 1 kHz and 10 Hz. The Csc values were nearly frequency independent. This phenomenon is contradictory to what was published earlier on iron oxides formed in borate buffer [20,26]. The Csc obtained at 1 kHz are presented in Table 1. The CH values in Table 1 were derived with the equation proposed by Hsu and Mansfeld [55]: nH −1

 ) CH = Q(ωm

Fig. 5. Equivalent circuit representing the oxide/electrolyte system used for EIS data modelling.

resistance for the charge transfer step Rct and the total double layer capacitance Cdl which is considered to be equal to CH in the used conditions [54]. The second Voight element represents the space charge region in the substrate, with the space charge capacitance, Csc , in parallel with the electrical resistance of the iron oxide layer, Rox [20]. To take into account the non-ideal behavior of the capacitive elements (CH and Csc ), constant phase elements (CPE) were introduced in the models in the first place. The resistive elements and the CPE-parameters Q and n were determined through a fitting procedure. The obtained values for the Bode plots shown in Fig. 4 are presented in Table 1. The presented CPE dispersion coefficients (n) were characteristic for all studied oxide layers. nsc was never lower than 0.99, whereas nH was in the 0.83–0.86 range. Since nsc was nearly 1, the structure of the formed space charge layers was very homogenous indicating that uniform oxide surfaces were prepared. The space charge

(2)

 is the frequency at which the imaginary In Eq. (2), ωm part of the measured impedance (of the concerned relaxation) reached a maximum. According to the Mott–Schottky concept, the total capacitance Ctot of the semiconductor/electrolyte system is given by following equation:   1 1 1 2 = + 2 Csc CH Ctot   1 2 KT E − E = 2 + − (3) FB εε0 qND A2 q CH

where CH is the capacitance of the Helmholtz double layer, A the exposed electrode area, ε the dielectric constant of the oxide layer, ε0 the permittivity in vacuum (8.85 × 10−14 F cm−1 ), ND the doping concentration, E the applied potential, EFB the flatband potential, q the electronic charge and kT/q is about 25 mV at room temperature. Using the capacitance data (CH and Csc ) from the impedance data fittings, ND can be estimated from the −2 slope of the linear part of the Ctot versus E curve in the relevant −2 versus potential range and EFB from the extrapolation of the Ctot −2 E curve to Ctot = 0. In the case of semiconductors (type: Si, Ge, . . .) characterized by a doping concentration around 1016 cm−3 in depletion mode (for n-type if E > EFB ), CH is much larger than Csc as a result of the relatively thick space charge layer (nm range or larger). If the semiconductor is heavily doped, leading to a small value of the space charge layer thickness, Csc will be

Table 1 Fitting results of EIS data obtained at different potentials on a steel sample treated for 8 min in air at 250 ◦ C E (V vs. ref)

Re ( cm2 )

Rct (k cm2 )

QH × 104 (sn −1 cm−2 )

nH

CH × 104 (F cm−2 )

Rox (M cm2 )

Qsc × 104 (sn −1 cm−2 )

nsc

Csc × 104 (F cm−2 )

−0.3 −0.1 0.2 0.7

97.7 97.5 96.3 96.4

23.5 50.7 51.9 41.6

1.08 0.75 0.71 0.70

0.87 0.86 0.86 0.85

1.25 0.92 0.89 0.87

0.026 2.03 6.15 7.24

0.74 0.20 0.12 0.09

0.997 0.999 0.996 0.999

0.74 0.20 0.12 0.09

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Fig. 6. Mott–Schottky plots performed in borate buffer of several thermal oxide layers formed in N2 and air.

˚ range). In these conditions, CH may of the same order as CH (A not be neglected [20]. From the impedance data fitting, it was found that the capacitance of the space charge layer for the studied thermal oxide layers was only approximately 4 to 7 times smaller than the Helmholtz capacitance in the potential range −2 where Ctot versus E behaved linearly. So, the Helmholtz capacitance could not be neglected in the Mott–Schottky approach. In Fig. 6, typical Mott–Schottky plots for several thermal oxide layers are shown. The plots were acquired in both sweep directions but only those acquired in positive sweep direction are presented. No significant hysteresis was observed meaning that a limited oxide modification took place. At potentials between −0.3 and 0.8 V, the plots have a positive slope. At potentials more positive than 0.8 V, the Mott–Schottky plots start to decrease. As can be seen in the polarisation curves of Fig. 2, this potential region corresponds to the oxygen evolution onset. In this potential region, no direct correlation between the semiconducting properties and defect concentration can be established due to the onset of an electronic current resulting from the descence of the Fermi level under the redox level for the oxygen reduction and the tunneling through the Schottky barrier [20]. The positive slope of the Mott–Schottky plot at lower potentials, in the passive region, leads to the conclusion that the thermal oxide layers formed in air and N2 behave like n-type semiconductors [20]. Possible donors in the oxide structure are mainly Fe2+ metal ions and oxygen vacancies [56]. In order to recognize whether a linear dependence exists over a certain potential range, and thus, to obtain a reliable estimation for the doping concentra−2 tion, the derivative of Ctot with respect to the applied potential is calculated. Then, the doping concentration is determined in the potential region in which the derivative dC−2 /dE is constant. Following this procedure, a limited linear region at potentials between −0.2 and 0.2 V can be observed for all oxide layers. Non-linear behavior at high potentials is attributed in literature to several reasons. First of all, two (or more) donor levels in the oxide structure can be present [46,56]. The decreasing slope at high potentials can mean that trivalent iron vacancies start to contribute to the total doping concentration [57]. Secondly, the bulk properties (near the surface) of the oxide layer may change

as a result of changing ND with polarisation time or with distance form the surface [20]. As mentioned before, hysteresis was limited, so the non-linear behavior is not induced by time effects or ion movement under the high electric fields of the space charge region. Probably, the outer atomic layers of the iron oxide differ slightly from the deeper structure due to air or due to N2 contact during oxide formation at high temperature. Comparing the different oxide layers, the slopes of the Mott–Schottky curves of the different oxides clearly change. For the oxide layers formed in N2 , the slopes are lowest and do not show any variation with time in oven (and with magnetite thickness). The independency of the total capacitance with oxide thickness proves again that these oxides behave like semiconductors and not like dielectric media. In case of the air formed oxides, the slope is much larger but decreases with time in oven and total oxide thickness. The slope of the linear part of the Mott–Schottky plot is affected by the doping concentration ND and the exposed surface area (Eq. (3)). In order to obtain reliable values for ND and EFB , the effective exposed surface area in contact with the electrolyte has to be known. For this reason, the oxide roughness and effective exposed surface area were determined by atomic force microscopy. 3.3. Determination of the oxide roughness and effective surface area In Fig. 7, the RMS roughness of the oxide surfaces as a function of the total oxide thickness is shown for both atmospheres. At each point of the curve, the corresponding heat treatment time is labeled. The surface roughness of a polished steel sample is added as initial value for the substrate roughness. It is clearly visible that in both atmospheres the oxide surface roughness increases with time and oxide thickness. Comparing oxide layers formed in air and N2 with similar thickness, the surfaces of the air formed oxides are rougher. In case of 100 nm thick oxide layers, the RMS roughness of the air formed oxide is 33 nm, facing 7 nm for the oxide formed in N2 . No correlation has been found between the dispersion coefficients of the CPEs,

Fig. 7. Variation of the RMS roughness of the oxide surfaces formed in N2 and air as a function of the total oxide thickness. At each point of the curve, the corresponding heat treatment time is labeled.

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Fig. 8. Variation of the relative surface area of the oxide surfaces formed in N2 and air as a function of the total oxide thickness. At each point of the curve, the corresponding heat treatment time is labeled. The relative surface area is the surface area, obtained by AFM, divided by the projected surface area.

mentioned in Section 3.2, and the oxide surface roughness in the studied roughness range. The images do not show any cracks in the oxide surface, even for the thickest oxide layers (not shown here). The obtained roughness values can be used to calculate the relative imaged surface area. Fig. 8 shows the relative surface area of the oxides as calculated with the Nanoscope software. The relative surface area Ar is expressed by following equation: Ar =

AAFM Ap

(4)

where AAFM is the effective surface area obtained with AFM taking the roughness into account and Ap is the projected surface area, which is 1 ␮m2 . The relative surface area for polished steel is added and is equal to 1. For the air formed oxide layers, the relative surface area increases significantly with thickness reaching a value of 1.3 for the 100 nm thick layer. This means that the effective surface area in contact with the electrolyte is 30% larger than the projected surface area. For the oxides formed in N2 , the relative surface area is much lower and does not exceed a value of 1.1. The relative surface area was used to calculate the effective surface area in contact with the electrolyte for the different iron oxide layers. For instance, taking into account that the projected exposed surface area for the EIS measurements was 0.95 cm2 (Section 2.3), the effective exposed surface area of the oxide layer formed after 4 h in air was 1.26 cm2 . 3.4. Electronic and semiconducting properties of the oxide layers In Table 1, it can be seen that CH is roughly constant in the −0.2–0.8 V potential region. The defect concentration ND and flatband potential EFB were derived from the linear part of the Mott–Schottky curve using Eq. (3) with CH being the average of the CH values obtained in the −0.2–0.8 V region. An overview of the results is shown in Fig. 9. A dielectric constant value of 10 was assumed [26,53]. It can be seen that oxide formation at

Fig. 9. Doping concentrations and flatband potentials of several thermal oxide layers formed in N2 (a) and air (b) as a function of the total oxide thickness. The labeled values are the corresponding heat treatment times.

250 ◦ C in air leads to about a five times lower doping concentration in the outer oxide scale than oxides formed in N2 . This is expected as the hematite layer on top of the oxide film contains a lower concentration of Fe2+ dopants [32]. The doping concentration of the thermal oxides seems to remain approximately at the same value with heat treatment time and total oxide thickness. The space charge layer thickness (lsc (E)) at a certain potential E can be calculated by following equation: lsc (E) =

εε0 A Csc (E)

(5)

where A is the effective surface area and Csc (E) is the space charge capacitance at a certain potential E. For the oxides formed in N2 and air, the space charge layer thickness at 0.8 V was, respectively, around 0.45 and 0.9 nm. So, in case of the air formed oxide layers, lsc never exceeded the thickness of the hematite layer. Thus, the obtained oxide properties (ND and EFB ) are typical for the upper hematite scale. Since the doping concentrations of the oxide layers do not vary with treatment time and oxide thickness, it can be concluded that oxides with a relatively high hematite thickness will have a relatively high electronic resistance compared to oxides without or with a relatively thin hematite layer. The flatband potentials of the oxides formed in N2 and air are, respectively, around −0.43 and −0.33 V ver-

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Fig. 10. Oxide resistance as a function of potential extracted from EIS data modelling for several thermal oxide layers formed in N2 and air.

sus Ag/AgCl. The results are in good agreement with previous work on synthetic bulk iron oxides, like magnetite and oxidised magnetite at 250 ◦ C [26]. Fig. 10 presents the oxide resistance at different potentials extracted from EIS data fitting for the thermal oxide layers formed in N2 and air. For all oxide layers, a sharp increase of oxide resistance with three orders of magnitude near the flatband potential is observed and can be explained by the accumulation of negative charges at potentials below EFB (‘accumulation mode’) and the extraction of electrons from the oxide surface at potentials above EFB (‘Depletion mode’) [20]. At potentials between 0.0 and 0.8 V, the oxide resistance is very high and seems to decrease slightly with increasing potential. Above 0.8 V, a sharp decrease in oxide resistance is noted due to oxygen evolution and tunneling [20]. Since at the flatband potential no band bending occurs, the flatband potential can be considered as a reference potential for the semiconductor electrode. Table 2 displays the oxide resistances at the flatband potential of the studied oxide films, together with the corresponding hematite thickness as received from ellipsometry. These resistances are obtained by interpolating the oxide resistance values at the nearest applied potentials above and below the flatband potential. For the oxides formed in N2 , only an approximation of the oxide resistance at flatband potential can be made since the flatband potential is beyond the measured potential range. For the oxides formed in N2 , the oxide resistance at the flatband potential is the lowest Table 2 Oxide resistance Rox at the flatband potential EFB of several thermal oxides formed in N2 and air Thermal oxide

␣-Fe2 O3 thickness (nm)

Rox at EFB (k cm2 )

In N2 , 8 min In N2 , 30 min In air, 5 min In air, 8 min In air, 30 min In air, 2 u In air, 4 u

– – 6.5 6.9 2.6 2.5 2.7

<1 <1 13 17 11 7 8

of all studied oxide films and is independent on the magnetite thickness. Nevertheless, the oxide resistance is quite high compared to pure magnetite [42], probably meaning that a very thin distorted layer (<1 nm) was formed on top of the magnetite scale. The resistance of the air formed oxides is much higher and seems to be directly correlated to the hematite thickness. Since a correlation between the concentration of ionic defects and electronic conductivity can be made, it can be assumed as a first approximation that the local electronic conductivity is proportional to the local amount of ionic defects. So, the main contribution to the resistance of an oxide film will come from the region where the conductivity is minimal. This means that outer hematite layer will determine the conductivity of the oxide film. The resistance is proportional to the hematite film thickness. The resistivity ρox of the oxide layers can be expressed by following equation: ρox =

Rox A l

(6)

where Rox is the oxide resistance, A the effective surface area and l is the hematite thickness. For an oxide layer formed in air after 8 min at 250 ◦ C, the resistivity is about 5 × 108  m. Nearly, the same value has been found by Senkevich for hematite layers formed at 500 ◦ C [42]. These results confirm the presence of the hematite layer and show its influence on the electronic and semiconducting properties of the thermal oxide films. 4. Conclusions Iron oxides formed in air at 250 ◦ C consist of an inner magnetite (Fe3 O4 ) scale covered by a hematite (␣-Fe2 O3 ) layer. In N2 , magnetite films are formed. By combining linear voltammetry and ellipsometry, it was shown that the hematite thickness is inversed proportional to the dissolution rate of the underlying steel substrate. Due to the presence of a hematite layer, the open circuit potential of the steel substrate shifts in the anodic direction. By electrochemical impedance spectroscopy combined with the Mott–Schottky approach, it was shown that the semiconducting properties of thermally formed iron oxide layers could be varied by changing the atmosphere for thermal treatment. The doping concentration or charge carrier density ND of air formed oxide films was about five times lower than the oxides formed in N2 . More negative flatband potentials EFB were obtained for the oxides formed in N2 . In both atmospheres, it was noted that the semiconducting properties (ND and EF ) were independent of heat treatment time or oxide thickness. Moreover, it was shown that the total oxide resistance was determined by the thickness of the outer hematite scale. It was noted that steel surfaces treated for short times in air had oxide layers with a relatively thick hematite layer compared to long time treatments. The highest electronic oxide resistance was noted for steel surfaces treated for 8 min. In the near future, adhesion and delamination of organic coatings on top of the studied thermal oxides will be investigated. It is expected that the difference in semiconducting behavior will result in different delamination rates of organic coatings.

J. Wielant et al. / Electrochimica Acta 52 (2007) 7617–7625

Acknowledgements Jan Wielant is financed by the Flemish Institute for the promotion of scientific and technological research in industry (IWT). His work is realised with the financial support of Arcelor Research Industry Gent-OCAS nv. This work is part of the project entitled ‘Nanometer scale electron spectroscopy of organic layers on oxide surfaces’, financed by the Research Foundation—Flanders (FWO). References [1] S. Cotugno, D. Larobina, G. Mensitieri, P. Musto, G. Ragosta, Polymer 42 (2001) 6431. [2] J. Lange, B. Nicolas, J. Galy, J.F. Gerard, Polymer 43 (2002) 5985. [3] C.I. Elsner, E. Cavalcanti, O. Ferraz, A.R. Di Sarli, Prog. Org. Coat. 48 (2003) 50. [4] S. Bistac, M.F. Vallat, J. Schultz, Int. J. Adhes. Adhes. 18 (1998) 365. [5] F. Deflorian, L. Fedrizzi, J. Adhes. Sci. 13 (5) (1999) 629. [6] J.J. Suay, M.T. Rodriguez, K.A. Razzaq, J.J. Carpio, J.J. Saura, Prog. Org. Coat. 46 (2003) 121. [7] K. Wapner, M. Stratmann, G. Grundmeier, Silicon Chem. 2 (2003) 235. [8] G. Grundmeier, M. Brettmann, P. Thiemann, Appl. Surf. Sci. 217 (2003) 223. [9] J. Marsh, L. Minel, M.G. Barthes-Labrousse, D. Gorse, Appl. Surf. Sci. 133 (1998) 270. [10] J. van den Brand, S. Van Gils, P.C.J. Beentjes, H. Terryn, J.H.W. de Wit, Appl. Surf. Sci. 235 (4) (2004) 465. [11] J. van den Brand, O. Blajiev, P.C.J. Beentjes, H. Terryn, J.H.W. de Wit, Langmuir 20 (15) (2004) 6308. [12] J.B. Bajat, V.B. Miskovic-stankovic, Z. Kacarevic-popovic, Prog. Org. Coat. 47 (2003) 49. [13] K. Wapner, G. Grundmeier, Adv. Eng. Mater. 6 (3) (2004) 163. [14] V. Barranco, J. Carpentier, G. Grundmeier, Electrochim. Acta 49 (2004) 1999. [15] A. Leng, H. Streckel, M. Stratmann, Corros. Sci. 41 (1999) 547. [16] A. Leng, H. Streckel, M. Stratmann, Corros. Sci. 41 (1999) 579. [17] A. Leng, H. Streckel, K. Hofmann, M. Stratmann, Corros. Sci. 41 (1999) 599. [18] M. Stratmann, A. Leng, W. Fiirbeth, H. Streckel, H. Gehmecker, K.H. Gro␤e-Brinkhaus, Prog. Org. Coat. 27 (1996) 261. [19] W. F¨urbeth, M. Stratmann, Prog. Org. Coat. 39 (2000) 23. [20] S.R. Morrison, Electrochemistry at Semiconductor and Oxidized Metal Electrodes, Plenum Press, New York, 1980. [21] S.J. Oh, D.C. Cook, H.E. Townsend, Hyperfine Interact. 112 (14) (1998) 59. [22] D. Bersani, P.P. Lottici, A. Montenero, J. Raman Spectrosc. 30 (1999) 355. [23] S. Peulon, H. Antony, L. Legrand, A. Chausse, Electrochim. Acta 49 (2004) 2891.

7625

[24] B.-H. Kim, J.-Y. Lee, Y.-H. Choa, M. Higuchi, N. Mizutani, Mater. Sci. Eng. B107 (2004) 289. [25] S. Jain, A.O. Adeyeye, S.Y. Chan, C.B. Boothroyd, J. Phys. D: Appl. Phys. 37 (2004) 2720. [26] M. B¨uchler, P. Schmuki, H. B¨ohni, T. Stenberg, T. Mantyla, J. Electrochem. Soc. 145 (2) (1998) 378. [27] J. Szczyrbowski, G. Bruer, M. Ruske, J. Bartella, J. Schroeder, A. Zmelty, Surf. Coat. Technol. 112 (1999) 261. [28] L.J. Oblonsky, T.M. Devine, Corros. Sci. 37 (1) (1995) 17. [29] S.S. Kulkarni, C.D. Lokhande, Mater. Chem. Phys. 82 (2003) 151. [30] M. Cohen, J. Electrochem. Soc. 121 (6) (1984) 191C. [31] L.J. Oblonsky, A.J. Davenport, M.P. Ryan, H.S. Isaacs, R.C. Newman, J. Electrochem. Soc. 144 (7) (1997) 2398. [32] R.M. Cornell, U. Schwertmann, The Iron Oxides, second ed., Wiley-VCH, 2003. [33] J.M. West, Basic Corrosion and Oxidation, Ellis Horwood, Chichester, 1980. [34] S. Birosca, D. Dingley, R.L. Higginson, J. Microsc. 213 (3) (2004) 235. [35] J.B. Wagner, K.R. Lawless, A.T. Gwathmey, Trans. Met. Soc. 221 (1961) 257. [36] U.R. Evans, The Corrosion and Oxidation of Metals Scientific Principles and Practical Applications, Arnold, London, 1971. [37] M. Seo, J.B. Lumsden, R.W. Staehle, Surf. Sci. 50 (1975) 541. [38] A.F. Wells, Structural Inorganic Chemistry, Clarendon Press, Oxford, 1984. [39] A.J. Davenport, L.J. Oblonsky, M.P. Ryan, M.F. Toney, J. Electrochem. Soc. 147 (6) (2000) 2162. [40] D.D. Macdonald, J. Electrochem. Soc. 139 (12) (1992) 3434. [41] T. Zimina, E. Oshe, V. Dubkov, P. Zimin, E. Zabrodskaya, J. Electrochem. Soc. 145 (7) (1998) 2236. [42] J.J. Senkevich, D.A. Jones, I. Chatterjee, Corros. Sci. 42 (2000) 201. [43] M. Da Cunha Belo, B. Rondot, C. Compere, M.F. Montemor, A.M.P. Simoes, M.G.S. Ferreira, Corros. Sci. 40 (2/3) (1998) 481. [44] N.E. Hakiki, M. Da Cunha Belo, A.M.P. Simoes, M.G.S. Ferreira, J. Electrochem. Soc. 145 (11) (1998) 3821. [45] E.M.A. Martini, I.L. Muller, Corros. Sci. 42 (2000) 443. [46] S. Modiano, C.S. Fugivara, A.V. Benedetti, Corros. Sci. 46 (2004) 529. [47] R. Babic, M. Metikos-Hutovic, Z. Pilic, Corrosion 59 (10) (2003) 890. [48] V. Goossens, J. Wielant, S. Van Gils, R. Finsy, H. Terryn, Surf. Interface Anal. 38 (2006) 489. [49] E.D. Palik, Optical Constants of Solids, Academic Press, San Diego, CA, 1998. [50] F. Mansfeld, S. Lin, Y.C. Chen, H. Shih, J. Electrochem. Soc. 135 (4) (1988) 906. [51] D. Rahner, Solid State Ionics 86–88 (Part II) (1996) 865. [52] M. Bojinov, G. Fabricius, T. Laitinen, K. Makela, T. Saario, G. Sundholm, Electrochim. Acta 45 (2000) 2029. [53] M. B¨uchler, P. Schmuki, H. B¨ohni, J. Electrochem. Soc. 145 (2) (1998) 609. [54] H. Gerischer, Electrochim. Acta 35 (11–12) (1990) 1677. [55] C.H. Hsu, F. Mansfeld, Corrosion 57 (2001) 747. [56] E. Sikora, D.D. MacDonald, J. Electrochem. Soc. 147 (11) (2000) 4087. [57] M. Bojinov, T. Laitinen, K. M¨akel¨a, T. Saario, J. Electrochem. Soc. 148 (6) (2001) B243.