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Electrochemical behaviour of thermally treated Cr-oxide films deposited on stainless steel
Corrosion Science 44 (2002) 451±465
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Electrochemical behaviour of thermally treated Cr-oxide ®lms deposited on stainless steel M.J. Carmezim a, A.M. Sim~ oes b,*, M.O. Figueiredo c, M. Da Cunha Belo b a
Mechanical Engineering Department, Instituto Polit ecnico de Set ubal, EST, Campus IPS, 2910 Set ubal, Portugal b Chemical Engineering Department, Instituto Superior T ecnico, Av. Rovisco Pais, 1049-001 Lisboa Codex, Portugal c CENIMAT and Materials Science Department, FCT, Univ. Nova de Lisboa, 2825 Monte de Caparica, Portugal Received 22 May 2000; accepted 27 March 2001
Abstract The semiconductive behaviour of amorphous and crystalline Cr-oxide ®lms deposited on stainless steels was studied by capacitance and photoelectrochemical measurements. The amorphous ®lms presented metallic behaviour. For the crystalline ®lms, bandgaps were determined for direct and indirect transitions. Doping densities evaluated from Mott±Schottky plots were below 1019 cm 3 for the ®lms treated at high temperature, and decreased with the treatment temperature. The crystalline ®lms were described as a thick layer of non-stoichiometric chromium oxide with a bi-polar character. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Cr-oxide ®lms; Electrochemical behaviour; Photoelectrochemistry; Mott±Schottky plots
1. Introduction The electronic properties are of crucial importance regarding the protective character of passive ®lms on metals that exhibit a semiconductor or insulator behaviour [1±6]. A number of studies can be found in the literature on the semiconductivity of passive iron [7±10], for which n-type semiconductivity has been proven. *
Corresponding author. Tel.: +35-121-841-7234; fax: +35-121-840-4589. E-mail address: [email protected] (A.M. SimoÄes).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 1 ) 0 0 0 7 6 - 2
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Conversely, for passive chromium, results in the literature suggest that, although the properties may change deeply with the conditions of ®lm formation and applied potential, the electronic behaviour tends to correspond to a p-type semiconductor [5,11,12]. Passivity of stainless steels is commonly attributed to the formation of a Cr-enriched oxide thin ®lm on the surface. However, an investigation of passive ®lms formed on stainless steels has shown that the electronic properties are complex, once observed semiconductivity is potential dependent as a consequence of both chromium and iron oxides that compose the ®lm [6,13]. Recent research on arti®cial passive ®lms has tried to correlate the corrosion resistance with the composition of mixed oxide ®lms (Fe,Cr)2 O3 [14] and Fe2 O3 ± Cr2 O3 ±NiO [15]. Studies on the electronic properties of sputtered Fe- and Cr-oxides have been made in an attempt to simulate the passive ®lm behaviour [16]. In a previous study using thick Cr-oxide ®lms on stainless steel produced by chemical conversion plus heat treatment, a strong dependence of the protective eciency on the nature of the oxide phases was found [17]. The spinel phase obtained after thermal processing displayed improved resistance to pitting, when compared to an intermediate corundum phase. The objective of the present work was to further investigate the relationship between composition, crystal structure and electronic behaviour of these conversion oxide ®lms. 2. Materials and methods Chemical conversion Cr-oxide ®lms were prepared by immersion in a chromic± sulphuric acidic solution at 70°C, on commercial sheets of stainless steel AISI 446 (Cr 26.5%; Mn 0.8%; C 0.2%; N 0.2%; Si 0.5%; P 0.030%; S 0.015% w/w), followed by isothermal treatment in order to promote crystallisation of the oxide. The isothermal treatment was made under controlled H2 /H2 O atmosphere with dew point 40°C, at normal pressure and at a temperature of 300°C, 700°C, 800°C or 950°C. Further details on ®lm preparation may be found elsewhere [18]. The thickness of the ®lms was 250 nm. The ®lms formed at ambient temperature and those treated at 300°C were amorphous, whereas all the others were crystalline. Identi®cation by Xray diraction (XRD) has revealed that the ®lms treated at 700°C presented a single a-phase (trigonal corundum structure), whereas those treated at 800°C revealed a cphase (cubic spinel structure) co-existing with the corundum structure. For ®lms treated at 950°C the c-phase was the only phase present [18] ± Table 1. Elemental depth composition of those ®lms was previously studied by secondary ion mass spectrometry (SIMS) [19]. The oxygen depth pro®les showed an inward penetration of this species upon heating, indicating that the thickness of the oxide ®lm increased with the treatment temperature. Iron was present in all situations while manganese was only detected in ®lms treated at 950°C. According to an X-ray absorption spectroscopy (XAS) study, a 3 valence state was observed for chromium in all the ®lms down to a depth of 4 nm, while manganese with 2 valence was only detected in ®lms treated at 950°C, following the formation of spinel phase [17].
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Table 1 Films structure at dierent heat-treatment temperatures Heat-treatment temperature °C
Film structure
Ambient temperature 300 700 800 950
Amorphous Amorphous a ac c
In the present work, an anodic passive ®lm was studied along with the ®lms described above. To produce this ®lm, a sample of AISI 446 was cathodically prepolarised for 5 min, and then the ®lm grown by anodic polarisation at 800 mV during 1 h, in a borate buer pH 8.4 (0.075 M Na2 B4 O7 10H2 O, 0.3 M H3 BO3 ) solution at ambient temperature. The solution was prepared with analytical grade chemicals and distilled water, and de-aerated with nitrogen prior and during the potential imposition. A cell containing a platinum counter electrode and a saturated calomel reference electrode (SCE) was used to produce the anodic ®lm. The same cell arrangement and solution was used in all the electrochemical measurements. Potentiodynamic curves were obtained by polarisation in the anodic direction between 0.7 and 1:0 eV, with a scan rate of 1 mV s 1 . Capacitance data were obtained from impedance measurements at 1 kHz using a 10 mV a.c. signal and a step rate of 50 mV every 10 s, in the cathodic direction, for a potential range 0.6 to 1.0 V/(SCE). Potentials outside this range were excluded due to the instability of the solution, with either the reduction of hydrogen or the evolution of oxygen at more cathodic or more anodic potentials, respectively. Measurements were made using a lock-in ampli®er (EG&G 5210) coupled to a potentiostat (EG&G 273 A). The donor or acceptor density Nq was determined from the Mott±Schottky equation that can be written as [20]: 1 1 1 1 2 kT U U 1 fb C 2 CH2 Csc2 CH2 ee0 qNq q where CH is the Helmholtz capacitance and Csc the space charge capacitance, e the dielectric constant of the passive ®lm, e0 the vacuum permittivity, q the elementary charge for electrons ( e for holes), k the Boltzmann constant, T the absolute temperature and Ufb the ¯atband potential. The photocurrent was obtained using a 150 W Xe arc lamp (Oriel 6254) as light source, a monochromator (Oriel 77200), a mechanical chopper and a photodetector (Oriel 7183). The lock-in technique was applied, where intensity and phase of photocurrent at dierent wavelengths were measured by means of the two-phase lock-in ampli®er coupled to the chopper and to the potentiostat. The chopper frequency was 19 Hz and spectra were measured between 250 and 750 nm in steps of 50 nm. Intensities were normalised with respect to the photon emission of the light source. The quantum eciency, g, de®ned as the ratio between the photocurrent Iph , and the incident photon ¯ux, U0 , of energy hm, is given by Gartners'model [21]:
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g
Iph hm Eg qAw U0 hm
n=2
2
where A is a constant, w the thickness of the space charge layer, Eg the optical bandgap energy. The exponent n depends on the type of transition between the valence band and the conduction band. For a crystalline semiconductor, a direct phototransition corresponds to n 1, and an indirect phototransition to n 4. Optical bandgap energy was determined by plotting ghmn=2 vs. hm and extrapolating to g 0. Auger electron spectroscopy (AES) analysis was performed using a 10 keV primary electron beam with a current density of 50 nA and an incidence angle of 30° relatively to the sample surface. Sputter pro®les were obtained by Ar ion etching, carried out at 10 5 Pa.
3. Results 3.1. AES study The Auger depth pro®les (Fig. 1) revealed oxygen enrichment in the ®lm treated at 950°C (spinel phase) with respect to the one treated at 800°C. The bulk of the spinel is depleted in iron, whereas enrichment in manganese with respect to the substrate is observed. The presence of this species can only result from the substrate diusion during the heat treatment, and is in agreement with the previous XAS analysis [17]. Apart from a small increase in the iron concentration at the ®lm/electrolyte interface in the spinel, the ®lms can be considered reasonably homogeneous in terms of elemental depth composition. 3.2. Potentiodynamic study The polarisation plots of the deposited oxide ®lms reveal a passive range between 0.5 and 0:4 V (SCE). The corrosion potential was 0.7 V for all the oxide ®lms and also for the AISI 446 alloy (Fig. 2). In the passive range, the rate of anodic dissolution is similar for the ®lms treated at 700°C or 800°C and for the anodic ®lm. For the ®lm treated at 950°C the passive current is lower by one order of magnitude, revealing a more protective character of this ®lm. Some dissolution of chromium from the oxide occurred at 0.7±0.8 V [5,11], whereas above that potential, oxidation of the solution was observed with oxygen evolution. 3.3. Capacitance measurements The general trend of the capacitance data (Fig. 3) reveals a common pattern as far as the in¯uence of the potential is concerned. Repetition of the capacitance measurements in a second scan gave an unaltered response, indicating that the
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Fig. 1. Auger depth concentration pro®les of ®lms treated at (a) 800°C and (b) 950°C.
®lms remained stable in spite of the cathodic polarisation. The absolute values of capacitance decrease when the temperature of the thermal treatment increases. The capacitance of ®lms formed at ambient temperature and at 300°C (amorphous ®lms) is practically identical. Below 0.3 V, a maximum in capacitance is observed leading to very high values, in the range of 100 lF cm 2 . This capacitance is at list one order of magnitude higher than the typical values of passive ®lms, and also higher than the capacitance of the Helmholtz layer. Large capacitance values have also been observed by other authors [3,22], and cannot be attributed to a space charge layer but rather to some other process, such as movement of ions across the ®lm [3] or a faradaic process occurring at the oxide/electrolyte interface [23]. For the anodic passive ®lm, the capacitance values are smaller, with a plateau of about 10 lF cm 2 in the anodic region. For the crystalline ®lms, the capacitance is always below that of the anodic oxide. In these ®lms, the capacitance is always below 5 lF cm 2 in the range of potentials studied (Fig. 3b). The plot of capacitance vs. potential shows a plateau between 0.1 and 0.3 V. The capacitance in this plateau decreases with the heat treatment temperature, going from 2.5 lF cm 2 in the ®lm
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Fig. 2. Polarisation plots for the anodic ®lm formed on AISI 446 and for ®lms treated at various temperatures.
Fig. 3. Capacitance vs. potential variation for (a) the anodic ®lm formed on AISI 446 and the amorphous conversion ®lms and (b) ®lms treated at 700°C, 800°C and 950°C.
treated at 700°C to 0.4 lF cm 2 in the ®lm treated at 950°C. Above 0.3 V the capacitance increases in the ®lm treated at 700°C whereas below 0.3 V a maximum was observed in all the ®lms.
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Fig. 4. Mott±Schottky plots for (a) the anodic ®lm formed on AISI 446, the amorphous conversion ®lms and the ®lm treated at 700°C, and (b) ®lms treated at 700°C, 800°C and 950°C. The plot of the ®lm treated at 700°C is repeated for comparison. Values aected by a factor of 10 when indicated.
When the capacitance is plotted as 1=C 2 vs. potential (Fig. 4), two linear regions are observed, above and below a ¯atband potential of about 0:5 V (SCE). The existence of those two straight lines suggests Mott±Schottky behaviour of n-type (positive slope) and p-type (negative slope). The p-type behaviour was observed only in the crystalline ®lms, i.e., the ®lms formed at temperatures above 700°C, while the amorphous ®lms, formed at ambient temperature and at 300°C show practically a horizontal line below 0.5 V. The minimum of 1=C 2 in the crystalline ®lms (Fig. 4b) is shifted from the potential axis, indicating that the capacitance measured at 0.5 V does not correspond to the Helmholtz double layer. The doping densities determined from the capacitance plots indicate 1021 cm 3 for the amorphous ®lms, corresponding to a very high, nearly metallic conductivity. For the ®lms formed at high temperatures, doping densities are below 1019 cm 3 . The acceptor density is slightly higher than the donor density and both densities decrease with the temperature of ®lm treatment (Fig. 5). 3.4. Photocurrent measurements The tests made on the amorphous ®lms have not revealed any measurable sign of photocurrent. This result was expected, given the high conductivity of those ®lms. In contrast, the crystalline ®lms gave a photocurrent response in a range of potentials. The normalised photocurrent spectra obtained at a ®xed potential of 0.2 V
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Fig. 5. In¯uence of the temperature of the thermal treatment on the doping densities of ®lms.
(Fig. 6) reveals two maxima, one corresponding to the gap and another for a sub-gap transition, i.e., at larger wavelengths. The transition above the gap has the highest intensity in the ®lm treated at 700°C, and decreases with the treatment temperature, being approximately ®ve times lower in the ®lm treated at 950°C. The sub-gap transition occurs at 2.0 eV (610 nm) and is also more intense in the ®lm formed at 700°C. In the ®lm treated at 950°C, this transition practically vanishes. This transition is possibly related to the presence of Fe(II) in the oxide, and agrees with the value of 1.9 eV observed for the bandgap of iron oxide anodically formed on pure
Fig. 6. Normalised photocurrent spectra for the ®lms treated at various temperatures.
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Fig. 7. Determination of optical bandgap energy, assuming direct transition. Temperature of ®lm formation indicated in the plots.
iron [7±9] and of 2.1±2.2 eV determined for Fe2 O3 in bulk crystalline forms (supposedly a-type structure) [9,24]. The decrease in the intensity of the sub-gap transition also agrees with the decrease of iron content in the ®lm. An identical transition was observed on sputter-deposited Fe2 O3 by Virtanen et al. [16], and on passive ®lms formed on Fe±Cr±Mo alloys by Hakiki et al. [25]. Using the G artner approach, the best ®t was obtained for a direct transition (Fig. 7). A value of 3.5 eV was found in the ®lms treated at 700°C and 800°C, whereas 4.2 eV was determined in the ®lm treated at 950°C. A reasonable ®t was also found assuming an indirect transition (Fig. 8), giving bandgap energies of 2.9 eV for all the crystalline ®lms. These results are in reasonable agreement with other authors, in spite of the scatter that can be found in the literature. Bandgaps of 3.5 and 3.3 eV have been determined [5,26] for Cr2 O3 , although higher values have been reported
Fig. 8. Determination of optical bandgap energy, assuming indirect transition. Temperature of ®lm formation indicated in the plots.
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Fig. 9. Photocurrent as a function of potential at 310 nm for the ®lms treated at (a) 700°C (b) 800°C and (c) at 290 nm for the ®lm treated at 950°C.
[12,27]. Sunseri et al. [5] have studied the behaviour of passive ®lms potentiodynamically grown on pure chromium in sulphate solutions and obtained values of bandgap between 2.4 and 2.95 eV, depending on the pH, whereas 3.55 eV was determined for an air-formed passive ®lm. From these results, it was concluded that the lower values of bandgap compared to that of the air-formed ®lm was due to the hydration of the oxide [5]. Virtanen et al. [16] have measured the bandgap energy of sputter-deposited Cr2 O3 and found values of 2.4 and 2.7 eV for ®lms formed in the absence and in presence of oxygen, respectively, and explained the dierence by a dierence in the stoichiometry of the oxides and by their defect structure. The potential variation of the photocurrent was determined for the three crystalline ®lms in the cathodic direction and is presented in Fig. 9. Due to the observed dierences in the bandgap energy, the photocurrents were measured at dierent wavelengths. Thus, for ®lms treated at 700°C and 800°C, the photocurrent was measured at k 310 nm, whereas in the 950°C ®lm the measurements were made at 290 nm. In every case, the measurements were made after stabilisation at each potential (2 min/point). For the 700°C and 800°C ®lms, a cathodic photocurrent was measured below 1.0 V for both ®lms. At 0:8 0:1 V an inversion of the photocurrent took place, revealed by an abrupt change of the phase angle and by the dierence in the trend of the jIph j vs. potential plot. In these ®lms the potential dependence of the photocurrent around the ¯atband potential is approximately linear, i.e. Iph const U
Ufb;opt
3
where Ufb;opt is the optical ¯atband potential and U the applied potential. At potentials above 0:3 V the photocurrent decreases, possibly due to a large degree of recombination. Signi®cant dierences were observed in the ®lm treated at 950°C (c-phase). For this ®lm, the photocurrent was signi®cant smaller than in the previous ®lms, in
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agreement with the photocurrent spectra presented above. On the other hand, the potential dependence of the photocurrent in the anodic region is approximately parabolic [28], as expected for crystalline semiconductors [21]: Iph const U
Ufb;opt
1=2
4
Below 1.0 V the signal became very small and no inversion of the signal was observed in this case. The determination of the ¯atband potential from the change in signal of the photoresponse can therefore be made only in the ®lms treated at 700°C and 800°C, and are in the range 0.7 to 0.9 V. These values are below the ¯atband potential from the Mott±Schottky plots, which were 0.5 V. For the ®lm treated at 950°C the optical ¯atband potential is situated nearly 1.0 V below the Mott±Schottky ¯atband potential. Since the ®lm formed at 800°C is constituted of a- and c-phase, and given the small photocurrent produced in ®lms with a single c-phase, the results suggest that the increase in temperature leads to a thinning of the a-phase layer of the ®lm, accompanied by a thickening of the c-phase layer. The way in which the phases develop is however not clear. They can either develop as separate layers, or as intergrown phases.
4. Discussion For pure a-Cr2 O3 , p-type semiconductivity is usually accepted in literature [26], although doping with Fe(II) may lead to n-type semiconductivity [29]. Normal spinels such as chromite FeCr2 O4 are p-type semiconductors [30]. The p-type semiconductivity is common in spinels and can be explained by the existence of cation vacancies due to an increase of 3 vs. 2 species. From a chemical point of view, the ®lms studied in this work may be considered as practically homogeneous, consisting of chromium oxides with some iron and manganese. However, the Mott±Schottky plots are characterised by two regions, with a positive and a negative slope, separated by an apparent ¯atband potential. Such behaviour is compatible with the existence of two layers of dierent semiconductivity. The existence of a bilayer ®lm was postulated by Nagayama and Cohen [31] for iron, whereas the co-existence of n- and p-behaviour in the capacitance response was reported by Hakiki et al. [13] for oxide ®lms on stainless steel. That capacitance response of stainless steels was explained by the presence of two layers, an outer iron-rich layer and an inner chromium-rich layer corresponding to n- and p-semiconductivity, respectively. For the conversion ®lms studied in this work, however, the capacitance behaviour is consistent with the existence of cation vacancies near the surface and a relative excess of metal ions in the inner part, as earlier suggested by Sato [32]. Photoconductivity results from excitation across the forbidden gap of the constitutive oxides, creating an electron-hole pair which separates under the action of the electrical ®elds
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present in the space charge region. These ®elds can be localised at both the oxide± electrolyte and the oxide±metal interfaces and also, if a n±p junction or heterojunction is formed, in the internal regions of the ®lm, in the respective space charge regions. The transitions observed in the photocurrent measurements can be associated with the iron and chromium oxides. The sub-gap transition at 2 eV, assumed to be due to excitation of one electron from iron, can be assigned to the 5 T2g (Fe2 ) ! 6 A1g (Fe3 ) transition, according to Goodenough [33]. The decrease in photocurrent intensity between the alpha and the gamma phase can correspond either to an decrease in the ratio Fe2 /Fe3 , as observed by Virtanen et al. [16], or only to a decrease of the iron content in the ®lm. The latter hypothesis is supported by SIMS data [19]. The indirect transition at 2.9 eV is in agreement with values published for non-stoichiometric chromium oxide. The greatest dierence of the results in the present work compared to published data is related with the existence of direct transitions, usually not referred in the literature, but probably associated with the crystalline character of the oxide. The transition at 3.5 eV may be associated with nearly stoichiometric a-Cr2 O3 [5]. The high value of 4.2 eV obtained for the c-phase is dicult to interpret, and can be due to the presence of manganese in this oxide. 4.1. Electronic band structure model for the ®lm The results of the present study can be explained by the existence of a bi-polar ®lm, with a p-type semiconducting behaviour at the outer part of the ®lm, and an ntype semiconducting behaviour at the internal part following to the model described by Carlsson et al. for a classical semiconductor [34]. This model assumes that two active junctions account for the measurements: a n/p junction inside the ®lm, and a p/electrolyte junction at the oxide±electrolyte interface. At positive bias, i.e., under anodic polarisation, the band bending at the p/electrolyte side is relatively small, whereas the band bending at the n/p junction increases (Fig. 10). Under these conditions, an anodic photocurrent is generated, and the p/electrolyte junction constitutes the main barrier to the ¯ow of holes towards the electrolyte. When a cathodic bias is applied, the band bending near the electrolyte increases whereas the band bending near the n/p junction becomes comparatively less important. Under these conditions, the controlling factor for this current is the transport of holes at the
Fig. 10. Band model showing the bias responsible for anodic photocurrent; CB, VB: borders of conduction band and valence band, respectively.
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internal junction. Due to this barrier, holes will accumulate in the p-region, the band bending over the n/p junction will decrease, and the cathodic photocurrent will ¯ow. However, if this barrier is very large, it may happen that a large cathodic polarisation is necessary before this photocurrent becomes measurable. This may be the case with the ®lm treated at 950°C (c-phase), for which a cathodic photocurrent could not be measured. The minimum in the Mott±Schottky plot observed at 0.5 V is shifted from the potential axis. This shift can also be explained by the model. For a circuit with capacitances in series, the measured capacitance can be described by: 1 1 1 1 C Cn=p Cp=el CH
5
Cn=p and Cp=el being the capacitance of the n/p junction and p/electrolyte junction, respectively and CH the capacitance of the Helmholtz double layer. For junctions under conditions of depletion, the capacitance becomes much smaller than that of the Helmholtz layer, and thus 1 1 1 C Cn=p Cp=el
6
Although for homogeneous semiconductors these capacitances do not in¯uence the measurements at the ¯atband potential, in these oxides there is always one of the two junctions contributing to the capacitance measurements, therefore not allowing the measurement of the Helmholtz layer. Thus the response is controlled by the n/p junction, above the ¯atband potential, by the p/electrolyte junction below that potential, and by the resultant of the two at the ¯atband potential. The ®lm can thus be seen as an n-type semiconductor/p-type semiconductor bipolar ®lm. Near the surface, metal vacancies play the role of acceptors whereas the metal interstitial and/or oxygen vacancies localised at the inner part play the role of electron donors. The outer Cr-oxide ®lm presents p-type semiconductivity (in both corundum and spinel type structures) whereas the inner Cr-oxide ®lm present n-type semiconductivity doing to excess of cations or oxygen vacancies. 4.2. Corrosion resistance Based on polarisation measurements in solutions containing Fe(CN)36 /Fe(CN)46 redox systems, the pioneer work of Bianchi et al. [35] established an interesting relationship between the susceptibility to pitting corrosion of type 304 stainless steel and the semiconducting properties of the oxide ®lms formed on this material by thermal treatment. Those authors demonstrated that oxygen-defective ®lms (with respect to usual oxide stoichiometry), which manifest n-type semiconductivity favour the nucleation of pitting corrosion. On the contrary, the oxide ®lms with an excess of oxygen are more protective against pitting attack [35,36]. The relationships between the semiconductive character of the oxide ®lms obtained by chemical conversion plus thermal treatment and the corresponding pitting resistance are more complex. The present investigation shows that the ®lms are
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characterised by the presence of electrical barriers developed at the ®lm±electrolyte interface and at the junction. In this work the ®lms treated at higher temperature presented a lower doping density. These dopants are associated with localized states, situated in the bandgap, which may act as sites for pitting initiation, as suggested by Schmuki et al. [37]. The less defective ®lm should in this case have the best resistance, which is consistent with the behaviour of the ®lm treated at 950°C. The mechanism of corrosion protection may thus be described by a model similar to the bi-polar structure proposed by Sato [32], with the outer layer being enriched in oxygen and the inner layer enriched in metal ions or with oxygen vacancies. Under these conditions, these layers would become cation selective and anion selective, respectively, and this would confer the ®lm a protective character. 5. Conclusions The electrochemical properties of chromium oxide ®lms formed by chemical conversion and thermal treatment on stainless steel AISI 446 were studied by electrochemical and photoelectrochemical techniques. Based on the results, the following conclusions can be withdrawn: (1) The AES analysis of the ®lms has revealed a chromium oxide ®lm, with some iron and manganese, practically homogeneous in depth. The capacitance behaviour can be described by the existence of n- and p-type semiconductivity, with the Mott± Schottky response being determined by two junctions, a n/p junction in the bulk and a p/electrolyte junction at the surface of the ®lm. (2) Bandgap energies were determined as 2.9 eV for indirect transition and 3.5 or 4.2 eV for direct transition. A sub-gap transition was also observed at 2 eV and associated with the presence of iron in the ®lm. (3) Doping densities decrease when the temperature of ®lm treatment increases, being below 1019 cm 3 for ®lms treated above 800°C. (4) The ®lms can be described as being made of a thick layer of chromium oxide with a bi-polar character, with cation vacancies near the surface and excess of cations in the inner part, which can be responsible for the observed n- and p-type semiconductivity. Acknowledgements The authors thank Dr. F atima Montemor (ICEMS/IST) for the Auger analysis. References [1] [2] [3] [4]
H. Gerischer, Corros. Sci. 29 (1989) 257. H. Gerischer, Corros. Sci. 31 (1990) 81. A. Di Paola, Electrochimica Acta 34 (1989) 203. A.M.P. Sim~ oes, M.G.S. Ferreira, B. Rondot, M. Da Cunha Belo, J. Electrochem. Soc. 137 (1990) 82.
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[5] C. Sunseri, S. Piazza, F. Di Quarto, J. Electrochem. Soc. 137 (1990) 2411. [6] N.E. Hakiki, M. Da Cunha Belo, A.M.P. Sim~ oes, M.G.S. Ferreira, J. Electrochem. Soc. 145 (1998) 3821. [7] P.C. Searson, R.M. Latanision, U. Stimming, J. Electrochem. Soc. 135 (1988) 1358. [8] L.M. Abrantes, L.M. Peter, J. Electroanal. Chem. 150 (1983) 593. [9] S.M. Wilhelm, K.S. Yun, L.W. Ballenger, N. Hackerman, J. Electrochem. Soc. 126 (1979) 419. [10] M. B uchler, P. Schmuki, H. B ohni, J. Electrochem. Soc. 145 (1998) 609. [11] M. Metikos-Hukovic, M. Ceraj-Ceric, J. Electrochem. Soc. 134 (1987) 2193. [12] E.W.A. Young, J.H. Gerretsen, J.H.W. de Wit, J. Electrochem. Soc. 134 (1987) 2257. [13] N.E. Hakiki, S. Boudin, B. Rondot, M. Da Cunha Belo, Corros. Sci. 37 (1995) 1809. [14] P. Schmuki, S. Virtanen, H.S. Isaacs, M.P. Ryan, A.J. Davenport, H. B ohni, T. Stenberg, J. Electrochem. Soc. 145 (1998) 791. [15] K. Sugimoto, M. Seto, S. Tanaka, N. Hara, J. Electrochem. Soc. 140 (1993) 1586. [16] S. Virtanen, P. Schmuki, H. B ohni, P. Vuoristo, T. Mantyla, J. Electrochem. Soc. 142 (1995) 3067. [17] M.J. Carmezim, M.O. Figueiredo, in: Y. Pauleau, P.B. Barna (Eds.), Protective Coatings and Thin Films, NATO ASI Series, Kluwer Academic Publishers, 1997, pp. 431±439. [18] M.J. Carmezim, F.G. Carvalho, M.O. Figueiredo, Thin Sol. Films 258 (1995) 194. [19] M.J. Carmezim, O.M.N.D. Teodoro, M.O. Figueiredo, A.M.C. Moutinho, in: M. Andritschky (Ed.), Advances in Materials and Processing Technologies'97, Univ. Minho, Portugal, 1997, pp. 326±328. [20] R. de Gryse, W.P. Gomes, F. Cardon, J. Vennik, J. Electrochem. Soc. 122 (1975) 711. [21] W.W. G artner, Phys. Rev. 116 (1959) 84. [22] U. Stimming, J.W. Schultze, Ber. Bunsen. Phys. Chem. 12 (1976) 1297. [23] E.C. Dutoit, R.L. Van Meirhaeghe, F. Cardon, W.P. Gomes, Ber. Bunsen. Phys. Chem. 12 (1975) 1206. [24] F. Di Quarto, C. Sunseri, S. Piazza, M.C. Romano, J. Phys. Chem. B 101 (1997) 2519. [25] N.E. Hakiki, M. Da Cunha Belo, J. Electrochem. Soc. 143 (1996) 3088. [26] P.A. Cox, Transition Oxide Metals, Oxford University Press, 1992. [27] F.A.Kr oger, Corrosion (Houston), NACE-6, 1983, p. 89. [28] M.A. Butler, J. Appl. Phys. 48 (1977) 1914. [29] W.B. Sharp, Corros. Sci. 10 (1970) 283. [30] B. Gillot, F. Jemmali, F. Chassagneux, C. Salvaing, A. Rousset, J. Solid State Chem. 45 (1982) 317. [31] M. Nagayama, M. Cohen, J. Electrochem. Soc. 109 (1962) 781. [32] N. Sato, Corros. Sci. 31 (1990) 1. [33] J.B. Goodenough, Les Oxydes des Metaux de Transition, Gauthiers-Villars, Paris, 1967. [34] P. Carlsson, K. Uosaki, B. Holmstr om, H. Kita, J. Electrochem. Soc. 136 (1989) 524. [35] G. Bianchi, A. Cerquetti, F. Mazza, S. Torchio, Corros. Sci. 12 (1972) 495. [36] M.Z. Yang, J.L. Luo, Q. Yang, L.J. Qiao, Z. Qin, P.R. Norton, J. Electrochem. Soc. 146 (1999) 2107. [37] P. Schmuki, H. B ohni, J. Electrochem. Soc. 139 (1992) 1908.
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Electrochemical behaviour of thermally treated Cr-oxide ®lms deposited on stainless steel M.J. Carmezim a, A.M. Sim~ oes b,*, M.O. Figueiredo c, M. Da Cunha Belo b a
Mechanical Engineering Department, Instituto Polit ecnico de Set ubal, EST, Campus IPS, 2910 Set ubal, Portugal b Chemical Engineering Department, Instituto Superior T ecnico, Av. Rovisco Pais, 1049-001 Lisboa Codex, Portugal c CENIMAT and Materials Science Department, FCT, Univ. Nova de Lisboa, 2825 Monte de Caparica, Portugal Received 22 May 2000; accepted 27 March 2001
Abstract The semiconductive behaviour of amorphous and crystalline Cr-oxide ®lms deposited on stainless steels was studied by capacitance and photoelectrochemical measurements. The amorphous ®lms presented metallic behaviour. For the crystalline ®lms, bandgaps were determined for direct and indirect transitions. Doping densities evaluated from Mott±Schottky plots were below 1019 cm 3 for the ®lms treated at high temperature, and decreased with the treatment temperature. The crystalline ®lms were described as a thick layer of non-stoichiometric chromium oxide with a bi-polar character. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Cr-oxide ®lms; Electrochemical behaviour; Photoelectrochemistry; Mott±Schottky plots
1. Introduction The electronic properties are of crucial importance regarding the protective character of passive ®lms on metals that exhibit a semiconductor or insulator behaviour [1±6]. A number of studies can be found in the literature on the semiconductivity of passive iron [7±10], for which n-type semiconductivity has been proven. *
Corresponding author. Tel.: +35-121-841-7234; fax: +35-121-840-4589. E-mail address: [email protected] (A.M. SimoÄes).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 1 ) 0 0 0 7 6 - 2
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Conversely, for passive chromium, results in the literature suggest that, although the properties may change deeply with the conditions of ®lm formation and applied potential, the electronic behaviour tends to correspond to a p-type semiconductor [5,11,12]. Passivity of stainless steels is commonly attributed to the formation of a Cr-enriched oxide thin ®lm on the surface. However, an investigation of passive ®lms formed on stainless steels has shown that the electronic properties are complex, once observed semiconductivity is potential dependent as a consequence of both chromium and iron oxides that compose the ®lm [6,13]. Recent research on arti®cial passive ®lms has tried to correlate the corrosion resistance with the composition of mixed oxide ®lms (Fe,Cr)2 O3 [14] and Fe2 O3 ± Cr2 O3 ±NiO [15]. Studies on the electronic properties of sputtered Fe- and Cr-oxides have been made in an attempt to simulate the passive ®lm behaviour [16]. In a previous study using thick Cr-oxide ®lms on stainless steel produced by chemical conversion plus heat treatment, a strong dependence of the protective eciency on the nature of the oxide phases was found [17]. The spinel phase obtained after thermal processing displayed improved resistance to pitting, when compared to an intermediate corundum phase. The objective of the present work was to further investigate the relationship between composition, crystal structure and electronic behaviour of these conversion oxide ®lms. 2. Materials and methods Chemical conversion Cr-oxide ®lms were prepared by immersion in a chromic± sulphuric acidic solution at 70°C, on commercial sheets of stainless steel AISI 446 (Cr 26.5%; Mn 0.8%; C 0.2%; N 0.2%; Si 0.5%; P 0.030%; S 0.015% w/w), followed by isothermal treatment in order to promote crystallisation of the oxide. The isothermal treatment was made under controlled H2 /H2 O atmosphere with dew point 40°C, at normal pressure and at a temperature of 300°C, 700°C, 800°C or 950°C. Further details on ®lm preparation may be found elsewhere [18]. The thickness of the ®lms was 250 nm. The ®lms formed at ambient temperature and those treated at 300°C were amorphous, whereas all the others were crystalline. Identi®cation by Xray diraction (XRD) has revealed that the ®lms treated at 700°C presented a single a-phase (trigonal corundum structure), whereas those treated at 800°C revealed a cphase (cubic spinel structure) co-existing with the corundum structure. For ®lms treated at 950°C the c-phase was the only phase present [18] ± Table 1. Elemental depth composition of those ®lms was previously studied by secondary ion mass spectrometry (SIMS) [19]. The oxygen depth pro®les showed an inward penetration of this species upon heating, indicating that the thickness of the oxide ®lm increased with the treatment temperature. Iron was present in all situations while manganese was only detected in ®lms treated at 950°C. According to an X-ray absorption spectroscopy (XAS) study, a 3 valence state was observed for chromium in all the ®lms down to a depth of 4 nm, while manganese with 2 valence was only detected in ®lms treated at 950°C, following the formation of spinel phase [17].
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Table 1 Films structure at dierent heat-treatment temperatures Heat-treatment temperature °C
Film structure
Ambient temperature 300 700 800 950
Amorphous Amorphous a ac c
In the present work, an anodic passive ®lm was studied along with the ®lms described above. To produce this ®lm, a sample of AISI 446 was cathodically prepolarised for 5 min, and then the ®lm grown by anodic polarisation at 800 mV during 1 h, in a borate buer pH 8.4 (0.075 M Na2 B4 O7 10H2 O, 0.3 M H3 BO3 ) solution at ambient temperature. The solution was prepared with analytical grade chemicals and distilled water, and de-aerated with nitrogen prior and during the potential imposition. A cell containing a platinum counter electrode and a saturated calomel reference electrode (SCE) was used to produce the anodic ®lm. The same cell arrangement and solution was used in all the electrochemical measurements. Potentiodynamic curves were obtained by polarisation in the anodic direction between 0.7 and 1:0 eV, with a scan rate of 1 mV s 1 . Capacitance data were obtained from impedance measurements at 1 kHz using a 10 mV a.c. signal and a step rate of 50 mV every 10 s, in the cathodic direction, for a potential range 0.6 to 1.0 V/(SCE). Potentials outside this range were excluded due to the instability of the solution, with either the reduction of hydrogen or the evolution of oxygen at more cathodic or more anodic potentials, respectively. Measurements were made using a lock-in ampli®er (EG&G 5210) coupled to a potentiostat (EG&G 273 A). The donor or acceptor density Nq was determined from the Mott±Schottky equation that can be written as [20]: 1 1 1 1 2 kT U U 1 fb C 2 CH2 Csc2 CH2 ee0 qNq q where CH is the Helmholtz capacitance and Csc the space charge capacitance, e the dielectric constant of the passive ®lm, e0 the vacuum permittivity, q the elementary charge for electrons ( e for holes), k the Boltzmann constant, T the absolute temperature and Ufb the ¯atband potential. The photocurrent was obtained using a 150 W Xe arc lamp (Oriel 6254) as light source, a monochromator (Oriel 77200), a mechanical chopper and a photodetector (Oriel 7183). The lock-in technique was applied, where intensity and phase of photocurrent at dierent wavelengths were measured by means of the two-phase lock-in ampli®er coupled to the chopper and to the potentiostat. The chopper frequency was 19 Hz and spectra were measured between 250 and 750 nm in steps of 50 nm. Intensities were normalised with respect to the photon emission of the light source. The quantum eciency, g, de®ned as the ratio between the photocurrent Iph , and the incident photon ¯ux, U0 , of energy hm, is given by Gartners'model [21]:
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g
Iph hm Eg qAw U0 hm
n=2
2
where A is a constant, w the thickness of the space charge layer, Eg the optical bandgap energy. The exponent n depends on the type of transition between the valence band and the conduction band. For a crystalline semiconductor, a direct phototransition corresponds to n 1, and an indirect phototransition to n 4. Optical bandgap energy was determined by plotting ghmn=2 vs. hm and extrapolating to g 0. Auger electron spectroscopy (AES) analysis was performed using a 10 keV primary electron beam with a current density of 50 nA and an incidence angle of 30° relatively to the sample surface. Sputter pro®les were obtained by Ar ion etching, carried out at 10 5 Pa.
3. Results 3.1. AES study The Auger depth pro®les (Fig. 1) revealed oxygen enrichment in the ®lm treated at 950°C (spinel phase) with respect to the one treated at 800°C. The bulk of the spinel is depleted in iron, whereas enrichment in manganese with respect to the substrate is observed. The presence of this species can only result from the substrate diusion during the heat treatment, and is in agreement with the previous XAS analysis [17]. Apart from a small increase in the iron concentration at the ®lm/electrolyte interface in the spinel, the ®lms can be considered reasonably homogeneous in terms of elemental depth composition. 3.2. Potentiodynamic study The polarisation plots of the deposited oxide ®lms reveal a passive range between 0.5 and 0:4 V (SCE). The corrosion potential was 0.7 V for all the oxide ®lms and also for the AISI 446 alloy (Fig. 2). In the passive range, the rate of anodic dissolution is similar for the ®lms treated at 700°C or 800°C and for the anodic ®lm. For the ®lm treated at 950°C the passive current is lower by one order of magnitude, revealing a more protective character of this ®lm. Some dissolution of chromium from the oxide occurred at 0.7±0.8 V [5,11], whereas above that potential, oxidation of the solution was observed with oxygen evolution. 3.3. Capacitance measurements The general trend of the capacitance data (Fig. 3) reveals a common pattern as far as the in¯uence of the potential is concerned. Repetition of the capacitance measurements in a second scan gave an unaltered response, indicating that the
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Fig. 1. Auger depth concentration pro®les of ®lms treated at (a) 800°C and (b) 950°C.
®lms remained stable in spite of the cathodic polarisation. The absolute values of capacitance decrease when the temperature of the thermal treatment increases. The capacitance of ®lms formed at ambient temperature and at 300°C (amorphous ®lms) is practically identical. Below 0.3 V, a maximum in capacitance is observed leading to very high values, in the range of 100 lF cm 2 . This capacitance is at list one order of magnitude higher than the typical values of passive ®lms, and also higher than the capacitance of the Helmholtz layer. Large capacitance values have also been observed by other authors [3,22], and cannot be attributed to a space charge layer but rather to some other process, such as movement of ions across the ®lm [3] or a faradaic process occurring at the oxide/electrolyte interface [23]. For the anodic passive ®lm, the capacitance values are smaller, with a plateau of about 10 lF cm 2 in the anodic region. For the crystalline ®lms, the capacitance is always below that of the anodic oxide. In these ®lms, the capacitance is always below 5 lF cm 2 in the range of potentials studied (Fig. 3b). The plot of capacitance vs. potential shows a plateau between 0.1 and 0.3 V. The capacitance in this plateau decreases with the heat treatment temperature, going from 2.5 lF cm 2 in the ®lm
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Fig. 2. Polarisation plots for the anodic ®lm formed on AISI 446 and for ®lms treated at various temperatures.
Fig. 3. Capacitance vs. potential variation for (a) the anodic ®lm formed on AISI 446 and the amorphous conversion ®lms and (b) ®lms treated at 700°C, 800°C and 950°C.
treated at 700°C to 0.4 lF cm 2 in the ®lm treated at 950°C. Above 0.3 V the capacitance increases in the ®lm treated at 700°C whereas below 0.3 V a maximum was observed in all the ®lms.
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Fig. 4. Mott±Schottky plots for (a) the anodic ®lm formed on AISI 446, the amorphous conversion ®lms and the ®lm treated at 700°C, and (b) ®lms treated at 700°C, 800°C and 950°C. The plot of the ®lm treated at 700°C is repeated for comparison. Values aected by a factor of 10 when indicated.
When the capacitance is plotted as 1=C 2 vs. potential (Fig. 4), two linear regions are observed, above and below a ¯atband potential of about 0:5 V (SCE). The existence of those two straight lines suggests Mott±Schottky behaviour of n-type (positive slope) and p-type (negative slope). The p-type behaviour was observed only in the crystalline ®lms, i.e., the ®lms formed at temperatures above 700°C, while the amorphous ®lms, formed at ambient temperature and at 300°C show practically a horizontal line below 0.5 V. The minimum of 1=C 2 in the crystalline ®lms (Fig. 4b) is shifted from the potential axis, indicating that the capacitance measured at 0.5 V does not correspond to the Helmholtz double layer. The doping densities determined from the capacitance plots indicate 1021 cm 3 for the amorphous ®lms, corresponding to a very high, nearly metallic conductivity. For the ®lms formed at high temperatures, doping densities are below 1019 cm 3 . The acceptor density is slightly higher than the donor density and both densities decrease with the temperature of ®lm treatment (Fig. 5). 3.4. Photocurrent measurements The tests made on the amorphous ®lms have not revealed any measurable sign of photocurrent. This result was expected, given the high conductivity of those ®lms. In contrast, the crystalline ®lms gave a photocurrent response in a range of potentials. The normalised photocurrent spectra obtained at a ®xed potential of 0.2 V
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Fig. 5. In¯uence of the temperature of the thermal treatment on the doping densities of ®lms.
(Fig. 6) reveals two maxima, one corresponding to the gap and another for a sub-gap transition, i.e., at larger wavelengths. The transition above the gap has the highest intensity in the ®lm treated at 700°C, and decreases with the treatment temperature, being approximately ®ve times lower in the ®lm treated at 950°C. The sub-gap transition occurs at 2.0 eV (610 nm) and is also more intense in the ®lm formed at 700°C. In the ®lm treated at 950°C, this transition practically vanishes. This transition is possibly related to the presence of Fe(II) in the oxide, and agrees with the value of 1.9 eV observed for the bandgap of iron oxide anodically formed on pure
Fig. 6. Normalised photocurrent spectra for the ®lms treated at various temperatures.
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Fig. 7. Determination of optical bandgap energy, assuming direct transition. Temperature of ®lm formation indicated in the plots.
iron [7±9] and of 2.1±2.2 eV determined for Fe2 O3 in bulk crystalline forms (supposedly a-type structure) [9,24]. The decrease in the intensity of the sub-gap transition also agrees with the decrease of iron content in the ®lm. An identical transition was observed on sputter-deposited Fe2 O3 by Virtanen et al. [16], and on passive ®lms formed on Fe±Cr±Mo alloys by Hakiki et al. [25]. Using the G artner approach, the best ®t was obtained for a direct transition (Fig. 7). A value of 3.5 eV was found in the ®lms treated at 700°C and 800°C, whereas 4.2 eV was determined in the ®lm treated at 950°C. A reasonable ®t was also found assuming an indirect transition (Fig. 8), giving bandgap energies of 2.9 eV for all the crystalline ®lms. These results are in reasonable agreement with other authors, in spite of the scatter that can be found in the literature. Bandgaps of 3.5 and 3.3 eV have been determined [5,26] for Cr2 O3 , although higher values have been reported
Fig. 8. Determination of optical bandgap energy, assuming indirect transition. Temperature of ®lm formation indicated in the plots.
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Fig. 9. Photocurrent as a function of potential at 310 nm for the ®lms treated at (a) 700°C (b) 800°C and (c) at 290 nm for the ®lm treated at 950°C.
[12,27]. Sunseri et al. [5] have studied the behaviour of passive ®lms potentiodynamically grown on pure chromium in sulphate solutions and obtained values of bandgap between 2.4 and 2.95 eV, depending on the pH, whereas 3.55 eV was determined for an air-formed passive ®lm. From these results, it was concluded that the lower values of bandgap compared to that of the air-formed ®lm was due to the hydration of the oxide [5]. Virtanen et al. [16] have measured the bandgap energy of sputter-deposited Cr2 O3 and found values of 2.4 and 2.7 eV for ®lms formed in the absence and in presence of oxygen, respectively, and explained the dierence by a dierence in the stoichiometry of the oxides and by their defect structure. The potential variation of the photocurrent was determined for the three crystalline ®lms in the cathodic direction and is presented in Fig. 9. Due to the observed dierences in the bandgap energy, the photocurrents were measured at dierent wavelengths. Thus, for ®lms treated at 700°C and 800°C, the photocurrent was measured at k 310 nm, whereas in the 950°C ®lm the measurements were made at 290 nm. In every case, the measurements were made after stabilisation at each potential (2 min/point). For the 700°C and 800°C ®lms, a cathodic photocurrent was measured below 1.0 V for both ®lms. At 0:8 0:1 V an inversion of the photocurrent took place, revealed by an abrupt change of the phase angle and by the dierence in the trend of the jIph j vs. potential plot. In these ®lms the potential dependence of the photocurrent around the ¯atband potential is approximately linear, i.e. Iph const U
Ufb;opt
3
where Ufb;opt is the optical ¯atband potential and U the applied potential. At potentials above 0:3 V the photocurrent decreases, possibly due to a large degree of recombination. Signi®cant dierences were observed in the ®lm treated at 950°C (c-phase). For this ®lm, the photocurrent was signi®cant smaller than in the previous ®lms, in
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agreement with the photocurrent spectra presented above. On the other hand, the potential dependence of the photocurrent in the anodic region is approximately parabolic [28], as expected for crystalline semiconductors [21]: Iph const U
Ufb;opt
1=2
4
Below 1.0 V the signal became very small and no inversion of the signal was observed in this case. The determination of the ¯atband potential from the change in signal of the photoresponse can therefore be made only in the ®lms treated at 700°C and 800°C, and are in the range 0.7 to 0.9 V. These values are below the ¯atband potential from the Mott±Schottky plots, which were 0.5 V. For the ®lm treated at 950°C the optical ¯atband potential is situated nearly 1.0 V below the Mott±Schottky ¯atband potential. Since the ®lm formed at 800°C is constituted of a- and c-phase, and given the small photocurrent produced in ®lms with a single c-phase, the results suggest that the increase in temperature leads to a thinning of the a-phase layer of the ®lm, accompanied by a thickening of the c-phase layer. The way in which the phases develop is however not clear. They can either develop as separate layers, or as intergrown phases.
4. Discussion For pure a-Cr2 O3 , p-type semiconductivity is usually accepted in literature [26], although doping with Fe(II) may lead to n-type semiconductivity [29]. Normal spinels such as chromite FeCr2 O4 are p-type semiconductors [30]. The p-type semiconductivity is common in spinels and can be explained by the existence of cation vacancies due to an increase of 3 vs. 2 species. From a chemical point of view, the ®lms studied in this work may be considered as practically homogeneous, consisting of chromium oxides with some iron and manganese. However, the Mott±Schottky plots are characterised by two regions, with a positive and a negative slope, separated by an apparent ¯atband potential. Such behaviour is compatible with the existence of two layers of dierent semiconductivity. The existence of a bilayer ®lm was postulated by Nagayama and Cohen [31] for iron, whereas the co-existence of n- and p-behaviour in the capacitance response was reported by Hakiki et al. [13] for oxide ®lms on stainless steel. That capacitance response of stainless steels was explained by the presence of two layers, an outer iron-rich layer and an inner chromium-rich layer corresponding to n- and p-semiconductivity, respectively. For the conversion ®lms studied in this work, however, the capacitance behaviour is consistent with the existence of cation vacancies near the surface and a relative excess of metal ions in the inner part, as earlier suggested by Sato [32]. Photoconductivity results from excitation across the forbidden gap of the constitutive oxides, creating an electron-hole pair which separates under the action of the electrical ®elds
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present in the space charge region. These ®elds can be localised at both the oxide± electrolyte and the oxide±metal interfaces and also, if a n±p junction or heterojunction is formed, in the internal regions of the ®lm, in the respective space charge regions. The transitions observed in the photocurrent measurements can be associated with the iron and chromium oxides. The sub-gap transition at 2 eV, assumed to be due to excitation of one electron from iron, can be assigned to the 5 T2g (Fe2 ) ! 6 A1g (Fe3 ) transition, according to Goodenough [33]. The decrease in photocurrent intensity between the alpha and the gamma phase can correspond either to an decrease in the ratio Fe2 /Fe3 , as observed by Virtanen et al. [16], or only to a decrease of the iron content in the ®lm. The latter hypothesis is supported by SIMS data [19]. The indirect transition at 2.9 eV is in agreement with values published for non-stoichiometric chromium oxide. The greatest dierence of the results in the present work compared to published data is related with the existence of direct transitions, usually not referred in the literature, but probably associated with the crystalline character of the oxide. The transition at 3.5 eV may be associated with nearly stoichiometric a-Cr2 O3 [5]. The high value of 4.2 eV obtained for the c-phase is dicult to interpret, and can be due to the presence of manganese in this oxide. 4.1. Electronic band structure model for the ®lm The results of the present study can be explained by the existence of a bi-polar ®lm, with a p-type semiconducting behaviour at the outer part of the ®lm, and an ntype semiconducting behaviour at the internal part following to the model described by Carlsson et al. for a classical semiconductor [34]. This model assumes that two active junctions account for the measurements: a n/p junction inside the ®lm, and a p/electrolyte junction at the oxide±electrolyte interface. At positive bias, i.e., under anodic polarisation, the band bending at the p/electrolyte side is relatively small, whereas the band bending at the n/p junction increases (Fig. 10). Under these conditions, an anodic photocurrent is generated, and the p/electrolyte junction constitutes the main barrier to the ¯ow of holes towards the electrolyte. When a cathodic bias is applied, the band bending near the electrolyte increases whereas the band bending near the n/p junction becomes comparatively less important. Under these conditions, the controlling factor for this current is the transport of holes at the
Fig. 10. Band model showing the bias responsible for anodic photocurrent; CB, VB: borders of conduction band and valence band, respectively.
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internal junction. Due to this barrier, holes will accumulate in the p-region, the band bending over the n/p junction will decrease, and the cathodic photocurrent will ¯ow. However, if this barrier is very large, it may happen that a large cathodic polarisation is necessary before this photocurrent becomes measurable. This may be the case with the ®lm treated at 950°C (c-phase), for which a cathodic photocurrent could not be measured. The minimum in the Mott±Schottky plot observed at 0.5 V is shifted from the potential axis. This shift can also be explained by the model. For a circuit with capacitances in series, the measured capacitance can be described by: 1 1 1 1 C Cn=p Cp=el CH
5
Cn=p and Cp=el being the capacitance of the n/p junction and p/electrolyte junction, respectively and CH the capacitance of the Helmholtz double layer. For junctions under conditions of depletion, the capacitance becomes much smaller than that of the Helmholtz layer, and thus 1 1 1 C Cn=p Cp=el
6
Although for homogeneous semiconductors these capacitances do not in¯uence the measurements at the ¯atband potential, in these oxides there is always one of the two junctions contributing to the capacitance measurements, therefore not allowing the measurement of the Helmholtz layer. Thus the response is controlled by the n/p junction, above the ¯atband potential, by the p/electrolyte junction below that potential, and by the resultant of the two at the ¯atband potential. The ®lm can thus be seen as an n-type semiconductor/p-type semiconductor bipolar ®lm. Near the surface, metal vacancies play the role of acceptors whereas the metal interstitial and/or oxygen vacancies localised at the inner part play the role of electron donors. The outer Cr-oxide ®lm presents p-type semiconductivity (in both corundum and spinel type structures) whereas the inner Cr-oxide ®lm present n-type semiconductivity doing to excess of cations or oxygen vacancies. 4.2. Corrosion resistance Based on polarisation measurements in solutions containing Fe(CN)36 /Fe(CN)46 redox systems, the pioneer work of Bianchi et al. [35] established an interesting relationship between the susceptibility to pitting corrosion of type 304 stainless steel and the semiconducting properties of the oxide ®lms formed on this material by thermal treatment. Those authors demonstrated that oxygen-defective ®lms (with respect to usual oxide stoichiometry), which manifest n-type semiconductivity favour the nucleation of pitting corrosion. On the contrary, the oxide ®lms with an excess of oxygen are more protective against pitting attack [35,36]. The relationships between the semiconductive character of the oxide ®lms obtained by chemical conversion plus thermal treatment and the corresponding pitting resistance are more complex. The present investigation shows that the ®lms are
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characterised by the presence of electrical barriers developed at the ®lm±electrolyte interface and at the junction. In this work the ®lms treated at higher temperature presented a lower doping density. These dopants are associated with localized states, situated in the bandgap, which may act as sites for pitting initiation, as suggested by Schmuki et al. [37]. The less defective ®lm should in this case have the best resistance, which is consistent with the behaviour of the ®lm treated at 950°C. The mechanism of corrosion protection may thus be described by a model similar to the bi-polar structure proposed by Sato [32], with the outer layer being enriched in oxygen and the inner layer enriched in metal ions or with oxygen vacancies. Under these conditions, these layers would become cation selective and anion selective, respectively, and this would confer the ®lm a protective character. 5. Conclusions The electrochemical properties of chromium oxide ®lms formed by chemical conversion and thermal treatment on stainless steel AISI 446 were studied by electrochemical and photoelectrochemical techniques. Based on the results, the following conclusions can be withdrawn: (1) The AES analysis of the ®lms has revealed a chromium oxide ®lm, with some iron and manganese, practically homogeneous in depth. The capacitance behaviour can be described by the existence of n- and p-type semiconductivity, with the Mott± Schottky response being determined by two junctions, a n/p junction in the bulk and a p/electrolyte junction at the surface of the ®lm. (2) Bandgap energies were determined as 2.9 eV for indirect transition and 3.5 or 4.2 eV for direct transition. A sub-gap transition was also observed at 2 eV and associated with the presence of iron in the ®lm. (3) Doping densities decrease when the temperature of ®lm treatment increases, being below 1019 cm 3 for ®lms treated above 800°C. (4) The ®lms can be described as being made of a thick layer of chromium oxide with a bi-polar character, with cation vacancies near the surface and excess of cations in the inner part, which can be responsible for the observed n- and p-type semiconductivity. Acknowledgements The authors thank Dr. F atima Montemor (ICEMS/IST) for the Auger analysis. References [1] [2] [3] [4]
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