The Composition of the Surface During Passivation of Stainless Steels*
INGEMAR OLEFJORD and BENGT-OLOF ELFSTROM*
Abstract
Molybdenum alloyed austenitic stainless steels were exposed in 0.1 M HCI + 0.4M NaCI and polarized in the active and passive regions of the alloys. The treated surfaces were analyzed by utilization of the electron spectroscopy for chemical analysis (ESCA) method. The reaction products formed on the surface after passivation consist mainly of chromium oxide and hydroxide. For the high Mo alloyed steel (6 wlo), even Mo is enriched in the passive film. Nickel is only present in the oxide to a very low extent. It is shown that polarization of the steels to their active potentials causes enrichment of the alloying elements—Cr, Ni, and Mo—in the outermost layers in the metal phase. Quantitative analysis indicates that the interphase, so formed, consists of only a few atomic layers. It is suggested that the interphase is formed by selective dissolution of iron during the initial period in the active state. The enrichment is con- sidered possible by a mechanism similar to short range ordering. It is postulated that the beneficial effect of nickel and molybdenum is not related to their occurrence in the passive film. Instead the collection of the alloying elements on the surface lowers the dissolution rate in the active phase and thereby provokes formation of the passive film during passivation and during repassivation in conditions of local attacks.
Introduction
The corrosion resistance of high alloyed steels is due to their ability to protect themselves by forming so-called passive films on their surfaces. The stability of the passive state is variable. It aggressive anions, such as halogens or oxidizing agents are present in the environment, the passive film is broken down and the steel is attacked locally. By alloying Cr steel with Ni and Mo, the corrosion properties are markedly im- proved. The influence of the alloying elements and their inter- action with the species in the liquid phase have been studied by electrochemical methods for many years. Review articles dealing with proposed models and phenomenological studies are given in the literature.1$
Since the advent of surface sensitive techniques such as electron spectroscopy for chemical analysis (ESCA) and Auger spectroscopy, it has been possible to analyze the reaction products formed on the surface.9.28 In our earlier ESCA studies,8-19 we have demonstrated that the passive film formed on stainless steel is enriched in chromium, whether it be ex- posed in sulfuric acid,13'16'18,23 in hydrochloric acid,18 or in neutral environments.12,"4'17'19 The molybdenum content of the passive film is not markedly different from its concentration in the alloy. It has also been clarified that the Ni content in the film is very low. Therefore, it has been suggested that the positive effects of the two alloying elements, Mo and Ni, do not arise from their forming oxidic products on the surface. We also have indications that during anodic dissolution, the alloy- ing elements—Cr, Ni, Mo—are enriched on the surface in their metallic states. All other studies were semiquantitative. The aim of this paper is to give a more quantitative description of the results.
* Presented during Corrosion/81 (Paper 245), April, 1981, Toron- to, Ontario.
*Department of Engineering Metals, Chalmers University of Technology, Goteborg, Sweden.
Equipment
The laboratory is equipped with an ESCA instrument (Hewlett Packard, model 5950 A) and a scanning Auger micro- probe (Physical Electronics, model 545). These two instru- ments are built together with a specimen preparation and handling system, which has been developed and built at this department. The setup is shown schematically in Figure la.
The main parts of the system are: an electrochemical cell for polarization of samples in aqueous solutions; a furnace for an- nealing in vacuum or in any desired gas up to 1000 C;29 a frac- ture device for fracturing metal samples at low temperature in ultra high vacuum;30 and two ion etching stations, in one of which the sample can be rotated in its own plane. The instru- ments, the manipulator chamber, and the furnace are separately ion pumped and they are connected to each other by valves. A transportation system "allows the sample to be moved between the different stations. The ESCA and the Auger analyzers are used independently of each other. The
"furnace" chamber is used as a sluice to introduce the samples into the system. Ten samples can be loaded in the same time.
The electrochemical cell (Figure 1b) is located inside a closed glass vessel, which is connected via a valve to the vacuum system. Before the sample is introduced into the cell and during the whole experiment, pure nitrogen gas is streamed through the vessel to empty it from oxygen. To avoid spontaneous passivation of the high alloyed steels, the sam- ple is polarized to a cathodic potential when the electrolyte is poured into the cell. Then the potential is swept in the anodic direction to the dissolution potential, where it is held until the surface is free from peroxide and its composition does not change with time (details will be given later). Starting from the corrosion potential, the potential is swept in one step to the desired potential either in the active or passive region. The polarization is interrupted by pouring acetone into the cel]. The current is switched off due to the low conductivity of acetone.
TABLE 1 — Chemical Composition of Test Materials in
Atomic Percent
No. Cr Ni Mo Fe
21.7 17.3 3.6 54
18.7 11.2 1.7 65
Note: Steels 1 and 2 are of commercial grade from Avesta AB and Sandvik AB, respectively. The alloys were exposed in 0.1 M HCI, 0.4M NaCI, and analyzed after polarization for 1 hour to the corrosion potential, active and passive potentials.
Before introducing the sample into the vacuum system, it is soaked in water free methanol. All exchange of liquids are per- formed by running the new liquid into the cell at the same rate at which the previous one is drawn off. Thereafter, the sample is transferred to the "furnace" chamber in pure nitrogen. All these precautions are taken to minimize the influence on the surface of oxygen in the atmosphere. For example, after treat- ment at the corrosion potential, it is possible to maintain a pure Ni surface in its metallic state,31 while more easily oxi- dized elements are oxidized to one or two oxidic layers during the transfer and pumping down of the chamber.
Material
As pointed out in the introduction, a series of stainless steels and pure metals has been studied after exposure in acids and neutral solutions. This paper intends to give a sum- mary of the results obtained at this department. The most systematic studies have been performed on pure Ni32 and the alloys FeNi,31 Fe-Cr-Mo16,18,23 after exposure in sulfuric acids.
Results from these earlier investigations will be mentioned, while for details of the experimental parts, reference is made to the original papers. However, this paper will be based on un- published results. The concentrations in atomic percent of the main alloy elements are given in Table 1. The concentrations of the minority elements are not given because they wilt not be discussed.
Quantitative Analysis
In the following, interest is focused on the analyses of thin oxide films and the metal underneath the passive film. It is possible to analyze the Jatter due to the fact that the thickness of the oxide is of the order of magnitude of the escape depth of the photoelectrons.
The integral intensity of the detected electrons which are created in a surface layer of thickness, a, can be expressed by the formula:
Mi = oJa X•DM•vMi•F•exp(—x/XMi)dx (1) where the indices, M and i, refer to the element and the atomic level respectively; X is the X-ray flux, D is the atomic density of the species, o is the differential photoionization cross section.
The factor F includes the product of the geometrical, the elec- tron optical, the analyzer transmission, and the detector effi- ciency functions; X is the attenuation length of the electrons.
The expression is discussed in detail elsewhere.33,34
The correct composition of the surface layer could be obtained from Equation (1) if all parameters were known. How- ever, this is not the case. The instrument response function F and the X-ray flux X change with time due to aging of the X-ray anode and the detector. Further, adequate cross section data do not exist. Therefore, to make a quantitative analysis, it is necessary to make calibration measurements to obtain the relative photoelectron yields from the elements which are of interest.
By integrating Equation (1) and using calibration con- stants, a simplified expression for the intensity obtained from a specific electron level of a certain M+ ion in an oxide is:
lox = a • Y DM. ox Coxw . m .
XM (1 — exp(— a/X )) (2)
The corresponding formula from the metal underneath is:
I met = a . Y Dmet Cmet smet exp(— alXox) (3)M M' M •M M M The a value is, in principle the product of the X and the F factors above. Y is the relative yield for the element M and is assumed to be the same in the oxide and the metal. The D values are the density of the cation, M+, or the element, M, in the calibration substances. The C-values are the atomic frac- tion of the species in the compound or alloy, which are in- vestigated. The two notations of the mean free path is made to mark that the X value differs in the metal and in the oxide phases. The FWHM and the background of the signals from the metallic and the oxide phases are different. At evaluation of the intensities from the spectra, this is taken into account by using apparent line widths, which give the same value for the experimental yield factor, YM, in the two states.
The exact value of the mean free path of the electrons, X, is not even known. From the literature,34 it appears that there is a great variation in the data. The same reference gives that the inelastic mean free path is about 1.2 nm for electrons with a kinetic energy corresponding to the 2p level of Cr in its metallic state.34 By using the old approximation that the at- tenuation length is proportional to the square root of the kinetic energy of the electrons, the corresponding values for Ni 2p, Fe 2p, and Mo 3d signals are: 1.0, 1.1, and 1.5 nm, respectively.
The attenuation lengths in the oxides are calculated from measurements of pure metals and oxides. The metals were cleaned by ion etching, analyzed, and then oxidized in oxygen in the furnace attached to the sample handling system. The oxides obtained are thicker than the attenuation length of the electrons but thin enough to exclude charging of the surface.
The attenuation lengths of the electrons emitted from the ca- tions Niox, Feox, Crox, and Moox are 1.3, 1.5, 1.6, and 1.9 nm, respectively. The values of the mean free paths used in Equa- tions (2) and (3) are those given previously, multiplied with the factor cos 51.5, 51.5° being the angle between the normai of the surface and the spectrometer axis.
To be able to calculate the concentrations, the thickness of the oxide layer has to be estimated. This is simply done by taking the ratio between Equations (2) and (3) and then solving for a. Then, the contents of the cation and the metallic species are obtained from Equations (2) and (3), respectively. The un- known factor a is eliminated by solving the equations with respect to a • C and then forming the following ratios:
CÁ?w = "MIN+ CN+ (4)
C _met C met/E Cmet (5)
M — M N N
where the summation is made over all elements.
It was stated in the introduction that some elements are enriched in the outermost layers in the metal phase. Therefore, it is expected that the content calculated from Equation (5) is an apparent concentration, CM p r}d that this value will differ from the chemistry of the alloy, CMe . In the following, the con- centration in the outermost metallic layer, CM, is given. This is calculated from a model of element distribution shown in Figure 2. There, it is assumed that all excess metal atoms are collected in just one plane underneath an oxide. The CM value
obtained is:
CS — CM p — CMet exp(— 2.5/XMet)
M 1 — exp (— 2.51 M ) (6)
FIGURE 1 — (a) The ESCA/SAM instrument, (b) the elec- trochemical ceil.
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FIGURE 2 — Assumed distribution of enriched elements in the metal phase.
Surf ace Analysis
Figure 3 shows the ESCA spectra recorded after pretreat- ment of the surface, anodic dissolution, and passivation at
-0.1 and +0.5 V (SCE). The sample chosen to illustrate the spectra is the 3.6 Mo alloy. The signals from Ni, Fe, Cr, and Mo represent both the metallic and the oxide states of the elements. The presence of the signals from the metal shows that the oxide products are thin enough to transmit photoelec- trons from the metal underneath the oxide.
In the Ni spectra, satellite signals appear beside the two described states. The strongest measurable Mo signal is the 3d state. This consists of two peaks-3d5,2 and 3d312—which are separated only by 3.5 eV. The doublet of the metallic state Is marked in the figure by the dotted lines. The 3d312 line overlaps with the 3d5l2 line of oxide states when present. It is, therefore, not easy to obtain the exact valence state of the molybdenum oxide. On the other hand, the contribution from the metal and the oxide is easily found because the shapes of the signals of the two states are identical.
The spectra recorded after pretreatment of the sample by grinding on emery paper are shown in row 1. It appears that the oxide consists mostly of Fe- and Cr-oxide. The two other main alloy elements—Ni and Mo—are only slightly oxidized. Later the quantitative analysis will show that the Fe and Cr concen- trations in the film correspond to that in the alloy. The oxygen signal exhibits double peaks corresponding to OH- (high binding energy) and 02- ((ow binding energy). This shows that at least the outermost layer of the oxide product consists of hydroxide. A systematic study of austenitic stainless steels that were treated in the same way is given elsewhere.14
Rows 2 to 4 show the spectra obtained after exposure of the steel (3.6 Mo) to the chloride solution at the potentials noted in the figure. The oxide products of the passivated samples contain the cations of Fe, Cr, and Mo. The nickel con- tent in the oxide is low. The dominating element in the passive film (rows 3 and 4) is chromium. Molybdenum exists in more
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FIGURE 3— ESCA spectra of the 21.7 Cr-17.3 Ni-3.6 Mo alloy after polishing in water and polarization: at the cor- rosion potential; -0.1 V (SCE); +0.5 V (SCE).
than one oxide state. As was pointed out earlier, it is not easy to elucidate the exact chemical state of Mo. Qualitatively, it is observed that its chemical state depends on the potential to which the sample was polarized. At a low potential, the film is a mixture of low valence Mo-oxide, while at a high potential (+0.5 V), molybdenum is oxidized to its six valence state.
It appears from the spectra (row 2) that oxide products are detected even after exposure of the sample at the corrosion potential. It is suggested that the oxide was formed during rinsing and handling of the sample after it was withdrawn from the cell. The surface oxide is extremely thin. Later, it will be shown that the thickness of the oxide is equivalent to two ca- tion layers. Because the cation density of the oxide is about half the density of the metal, the detected cations cor- responds to one atomic plane of the metal. In other words, dur- ing handling of the surface, one atomic layer is oxidized.
The information obtained by the ESCA analysis of the two steels was quantified and the results are expressed in Figures 4 and 5. The surfaces were either polarized to the active or to the passive regions. It was ascertained that the potential did not reach the potential where pitting corrosion starts. Both alloys were activated by polarization 200 mV below the corro- sion potential.
From Figure 4a, it appears that the oxide products formed during handling of the sample after exposure to the corrosion potential consist of equal amounts of Cr and Mo. Ni and Fe are present only to a very low extent. Figure 4b shows that the ap- parent composition of Ni in its metallic state is 27 alo. This is about 10% higher than the Ni content of the alloy. Thus, Ni is enriched on the surface in the metal phase. The evaluation also indicates that Mo is enriched in the metal phase. The Cr content is a few percent units lower than that quoted in the chemical analysis of the steel.
Figure 4c shows that the thickness of the oxide product formed during handling of the samples that had been treated at the corrosion potential, is about 0.6 nm. As pointed out above, the cations in the oxide correspond approximately to one atomic plane of the metal. The composition of the oxide would represent the outermost crystal plane during exposure in the acid solution, if it is assumed that dissolution does not
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FIGURE 4 — Quantitative analysis of sample No. 1 (21.7 Cr-17.3 Ni-3.6 Mo): A. The concentration of the cations in the oxide; B. The apparent concentration in the metal phase; C. The thickness of the passiva film; D. The com- position of the outermost metal layers. The solid lines in Figures B and D show the chemical analysis of the alloys.
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FIGURE 5 — Quantitative analysis of sample No. 2 (18.7 Cr-11.2 Ni-1.7 Mo): A. The concentration of the cations in the oxide; B. The apparent concentration in the metal phase; C. The thickness of the passiva film; D. The com- position of the outermost metal layers. The solid lines in Figures B and D show the chemical analysis of the alloys.
take place during the rinsing operation. However, the apparent enrichment of Ni indicates that the depth of the region, within which the composition of the alloy changes, is more than one atom layer thick. The simplest model that can be applied is that the enriched surface zone during active dissolution con- sists of two planes; namely, one consisting of metal atoms in the oxide and the other of the outermost layer in the metal phase immediately under the oxide. The contents of the metal species in the latter were calculated by using the simple model in Figure 2 and Equation (6). The estimated value for Ni is 44%.
The average value over the two atomic layers is half of that.
Figure 4d shows the estimated concentration in the outermost region for the different elements. It appears that the contents of all are in the range 20 to 30%. Thus, the Fe concentration is lowered by preferential dissolution and the alloy elements are enriched on the surface. Chromium and Ni are each enriched by a factor of 1.5, while Mo is enriched by a factor of 7.
The passive film formed on the surface after polarization to the passive region consists mainly of Cr-oxide. Figure 4a shows that the Cr content is about 50% at –100 mV and in- creases with the potential. Molybdenum behaves differently;
at the lower potential, the Mo content is rather high but falls under 10% at the higher potential. Nickel is only present to a few percent units while the Fe content is about 20%. Figure 4c shows that the thickness of the oxidic products increases with the potential reaching about 1.5 nm at the highest potential.
Equation (6) was also used for calculation of the composi- tion of the top layers in the metal phase underneath the passive film, after polarization to the passive potentials (Figure 4d). The validity of this is perhaps debatable because the passive film grows in the solution. However, this approach indicates as expected that the amounts of Fe and Ni in the top layers are about the same as in the active state and that Cr is depleted due to enrichment in the passive film. This demon- strates that Fe is selectively dissolved even during passiva- tion, otherwise Fe atoms would be collected underneath the passive film.
The Ni and the Fe contents in the passive film formed on the low Mo alloyed steel are about the same as detected on
POLISHED ON EMERY PERS:
A) In water B) In methanol
100 COMP. IN THE 0UTERMOST
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FIGURE 6 — The composition of the oxide and the ap- parent metal concentration of sample No. 1(21.7 Cr-17.3 Nl-3.6 Mo) after polishing In: A. water, B. methanol.
the high Mo alloyed sample (Figure 5a). The Mo content in the former is lower than in the latter. The maximum cation content of Cr is about 80%. At +200 mV, which is near the pitting potential, the Fe content increases at the same time as the Cr content decreases. This is even more pronounced at higher potentials (not shown in the figure).
In the active potential range, Fe is preferentially dissolved and the alloy elements are enriched in the outermost atomic layers. Figure 5d shows that at the corrosion potential, at least, the Fe content is almost the same and the enrichment of Ni and Mo is as high as for the high Mo alloyed steel.
The estimated thickness of the passive film is about 1.0 nm after passivation at –200 mV (Figure 5c). At the highest potential in the true passive state, the thickness is about 1.5 nm.
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detailed description of these would be too involved for the present paper. In order to confirm that the handling of the samples following the electrochemical treatment does not in- fluence the resuits, we refer to the resuits in Figure 6. These were obtained from polishing the high alloyed Mo sample on emery paper in either water or methanol. Polishing in water creates a fresh metal surface which is immediately covered by corrosion products. The composition of the products is markedly different from that found after dissolution in acid.
Excluding Ni, the reaction products consist of cations cor- responding to their concentration In the steel. It is, therefore, suggested that the dissolution is very low and that the atoms available in the surface region form the oxide. However, upon prolonged exposure, the Cr content in the oxide increases simultaneously with a falling Fe content.12'14
The measured concentration of the elements underneath the oxide is closer to the chemical analysis of the alloy than that obtained after exposure in acid. This is logical, since very little change of the surface composition is to be expected. The difference between the measured composition and the compo- sition of the alloy is within the experimental error.
The thickness of the oxide product formed in water is 1.8 nm. The same thickness of the oxide is obtained when an ion etched surface of the same sample is exposed to the protect- ing nitrogen gas used during transportation of the sample to the vacuum chamber. The measured thickness of the reaction product formed on the surface after polishing in methanol was about 0.8 nm. This shows that adsorbed methanol inhibits ox- idation during transportation to some extent. The only element which is enriched on the methanol treated surface is Cr in its oxide state. Iron and Mo are present in the oxide in the same proportion as in the alloy. In the metallic phase, the measured composition of these elements agrees with the composition of the alloy. The increased Ni content in the metal is related to the Jack of Ni in the oxide product.
It can be concluded from Figure 6 that the influence of the methanol and the protecting gas during movement of the sam- ple from the celi to the analyzer can be neglected. The only ele- ment which is slightly enriched is Cr and its concentration in the oxide is far below that obtained in the true passive film and in the product formed after treatment at the corrosion poten- tial.
Electrochemical Studies
The changes of the surface composition may influence the electrochemical behavior. To test this, cathodic and anodic polarization studies were made in the test solution. To eliminate concentration gradients in the neighborhood of the surface, the rotating disc method was used.35 All experiments started 200 mV below the corrosion potential by dipping the freshly polished sample into the solution. Thereafter, the potential was swept in the anodic direction with a rate of 1 mV persecond.
Figure 7 shows polarization diagrams of the pure ele- ments—Cr, Fe, Ni and Mo—and of three alloys. The curves denoted 1 and 2 represent the alloys discussed above, while number 3 is recorded from an ordinary 1818 austenitic stain- less steel which contains less than 0.15 w/o Mo. It appears (Figure 7a) that Fe and Ni are not passivated at all in the actual solution (0.1M HCI, 0.4M NaCI). Chromium is passive over a broad range. Molybdenum behaves differently compared to the other elements. It does not show any characteristic active- passive transition and the anodic current does not increase steeply with the potential as for Fe and Ni. Instead, the current is extremely low in a 400 mV wide range. The current is of the same order of magnitude as the passive current of Cr. Above +150 mV, the anodic current increases markedly, ESCA analysis36 of a Mo sample that was polarized in the low anodic current range showed a very small contribution from a low valence oxide state of Mo. The coverage was suggested to be about one or even less than a monolayer. Hitherto, it is not clear whether the oxide state is present on the metal surface in the solution or if it is formed during handling of the sample.
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FIGURE 7 — Polarizatlon diagrams: A. The metals—Fe, Cr, Ni, and Mo; B. The alloys-1. (21.7 Cr-17.3 Ni•3.6 Mo);
2. (18.7 Cr•11.2 NI-1.7 Mo); 3. (18 Cr-8 Ni-0.1 Mo); C.
Enlarged scale of the cathodic and anodic diagrams around the corrosion potential of the metals and the alloys .
After polarization at +500 mV, voluminous products of six valence Mo-oxide is formed.
The polarization curves of the alloys show that the high Mo alloyed steel, number 1, is passive in a broad potential range. The other two alloys are sensitive to pitting corrosion above a critical potential which limits their passive ranges.
The critical current density is about the same for passivation of the two Mo containing steels, while it is an order of mag- nitude higher for the non-Mo alloyed steel.
Parts of the cathodic and anodic polarization curves are shown in enlarged scale in Figure 7c. The diagram of the third alloy is not shown because it coincides with the diagram of Mo. The following explanation of the electrochemical behavior will only be qualitative. The three quantities which are of in- terest are: corrosion current, corrosion potential, and Tafel slope of the hydrogen reaction. The corrosion current can, in principle, be obtained from the intercept of the extrapolated lines of the linear parts of the cathodic and anodic curves.
Because no pronounced Tafel slopes are obtained from the anodic reaction, it is not possible to get any exact value.
However, from the figure, it appears that the self corrosion rate of Cr is about one order of magnitude higher than for Fe, which is another order of magnitude higher than for Ni and Mo.
The corrosion current of Mo is somewhat lower than for Ni. No marked difference between the corrosion currents of the two alloys can be noticed. The corrosion currents seem to be about the same as for Ni. Nickel shows the highest corrosion poten- tial. The corresponding potential of Mo is 40 mV lower. The two Mo alloyed steels have about the same corrosion potential as Ni. Figure 7b shows that the corrosion potential of the 18/8 alloy is markedly lower than for the others. The Tafel slope of the hydrogen evolution on Ni is about 100 mV per decade. Of all the slopes, this value is closest to the theoretical value, 120 mV/decade, which is obtained if proton discharge is the rate determining step. The slope of the cathodic reaction on Mo is markedly steeper than on pure Ni.
Discussion
This study confirms earlier investigations that the main compounds in the passive film formed on stainless steels are oxide products of Cr. The same resuits are obtained for any steel exposed to sulfuric-hydrochloric acids,16'18220,22,23,37
neutral solution,17'19 or to pure water for a long time.12,14,24 The Fe content in the film is a few tens of percent units. Nickel is only oxidized to a very low extent. It is therefore suggested that the positive influence of Ni on the corrosion properties of the steels is not its occurrence in the passive film. Instead, it will be argued that it is the enrichment of Ni at the.surface in the underlying metal phase which provokes the passivibility of the alloy. The role of Mo seems to be even more complex. In earlier reports where low Mo alloys were studied, Mo was not enriched in the oxide. The results discussed above show more clearly that Mo is also significantly enriched in the oxide prod- ucts. On the other hand, they also show that the Mo enrich- ment is high at low potential and decreases when the potential is increased. From the spectra, it appeared that the polariza- tion potential influences the valence state of Mo. At the highest potential, Mo is enriched to its six valence state. This behavior is similar to that of pure Mo. Results36 indicate that the reaction product consists of a very thin layer of Mo-oxide when Mo is in a low valence state, while at high potentials, voluminous six valence oxide is formed. However, the high Mo alloyed steel is passive up to a higher potential than that at which strong oxidation of Mo starts. Even when the Mo con- tent is lowered at high potential by oxidation and dissolution, Mo seems to have a beneficial effect on the stability of the passive state. One effect, suggested by Sugimoto and Sawda20 could be that the passive film (chromium oxyhydrox- ide) is stabilized by Moe. They also claim that the thickness of the passive film increases with the Mo content of the alloy.
This effect is not observed in our results. The film thickness is only dependent on the potential and not on the Mo content in the alloy. This potential dependence of the thickness has also been observed in ferritic steels exposed in sulfuric acids.18 In that case, the agreement between our ESCA results and ellip- sometry measurements38 was excellent.
The ESCA analyses show that during active dissolution, the alloying elements Cr, Ni, and Mo are enriched on the sur- face in their metallic state. An interphase, the composition of which differs from the chemical analysis of the alloy, is there- by created. It has been postulated18 that the chemical com- position of this interphase controls the corrosion properties of the alloy. The general basis for this statement is that electro- chemical reactions take place over an interface between the metal and the solution. The reaction rate depends on which metal the reaction occurs and the composition of the solution;
or more precisely the composition of the Helmholtz double layer. For an alloy, the electrochemical behavior would change if the composition of the surface changes over a few atomic layers.
It is proposed that the enrichment of the alloying elements occurs during the initial period during which iron is selectively dissolved. Subsequently, a stationary state is reached where the dissolution of each alloying element cor- responds to its composition in the alloy. The polarization studies of pure metals show that the dissolution rates of Ni and Mo are lower than that of Fe. This may explain why Ni and Mo are enriched on the surface. The case of Cr is more dif- ficult. Pure Cr is dissolved at a higher rate than pure Fe and yet it also seems to be enriched in the outermost interphase. This paradox has been explained with the words: synergistic effect between the species. If it is accepted that an interphase is formed on the surface, this phase is necessarily thermo- dynamically more stable than a phase of the alloy composi- tion. Because the surface phase is very thin, it must also be assumed that no new interface is created; instead, the two phases must be coherent so that the interfacial energy be- tween the two metal phases can be neglected.
A characteristic of the quarternary alloy—Fe, Cr, Ni,Mo- is that it forms intermetallic phases (a, X, Laves). These are
normally formed at high temperatures. The driving force to form these phases from a solution is that the attractive inter- action between unlike atoms is stronger than the bonding be- tween like atoms. It is suggested that during anodic dissolu- tion of the alloy, a similar interaction occurs; the solvent, Fe, is dissolved and interaction occurs between the remaining alloy- ing elements to form a more stable interphase on the surface.
The measured compositions of the interphases both for the high and low Mo alloyed steels are within the existence range for the a-phase.39 Of course, it cannot be claimed that the structure of a-phase is obtained on the surface. Instead, it is postulated that a mechanism akin to short range ordering oc- curs on the surface during anodic dissolution.
The "total" electrochemical behavior of a metal is a synergistic effect between the chemistry of the surface of the metal and the conditions in the Helmholtz double layer. To test if the composition of the metal phase influences the total hydrogen reaction, cathodic polarization curves were re- corded. These show as expected that there is no simple cor- relation between the overpotentials of the alloys and the pure metals. But they also show that the corrosion potential of the alloys is close to that of the more noble elernents. This is a fur- ther indication that the composition of the surface changes.
This effect has been demonstrated40 more effectively on an or- dinary 1818 alloy (alloy 3 in Figure 7, Mo content 0.15 wlo), which was exposed in 2.5M H2SO4. When cycling the potential between a cathodic and an anodic potential a number of times, the corrosion potentials was shifted in the anodic direc- tion during the second cycle. Surface analysis revealed that also in this case Mo and Ni were enriched. In another study31
of a single crystal of Fe 75 Ni, it was shown that in this case, the outermost atomic layer is almost completely covered by Ni atoms and that the corrosion potential agrees with the corro- sion potential of pure Ni.
The new experimental data presented in this paper con- firms the earlier results that the alloy elements are enriched on the surface in their metallic states during active dissolution. It is suggested that the enrichment provokes passivation by lowering the dissolution rate in the active phase preceding passivation. The high concentration of the passive film form- ing elements in the surface region reinforces the passivibility of the alloy. It has also been suggested18 that it is not only the general corrosion behavior of stainless steels that is influ- enced by the enrichment mechanism. The increased contents of some alloy elements in an interphase might also improve the resistance to pitting and crevice corrosion. The repassiva- tion of an initiated pit might be facilitated by enrichment of the interphase-stabilizing elements.
To get a complete view of the corrosion mechanisms of such a complicated alloy system as stainless steels, it is necessary to make systematic surface and electrochemical analyses on single crystals of the actual alloy system.
Conclusion
It is proposed that the enrichment of the alloy elements in the metal phase provokes passivation by lowering the dissolu- tion rate in the active phase preceding passivation; the alloy elements that form the passive film are availabie on the sur- face in enhanced concentration.
Acknowledgments
The authors are grateful to U. Jelvestam and G. Strom for their competent assistance in surface analysis and electro- chemical measurements, respectively. Financial support from the Swedish Board for Technical Development (STU) is grate- fully acknowledged.
References
1. Kolotyrkin, Ya. M. Corrosion, NACE, Vol. 19, p. 261t (1963).
2. Szklarska-Smialowska, Z. Corrosion, NACE, Vol. 27,p. 22 (1971).
3. Rosenfeld, 1. L. Proceedings of the Fifth International Congress on Metallic Corrosion, NACE, Vol. 53 (1974).
4. Szklarska-Smialowska, Z. Proceedings of the Fifth Inter- national Congress on Metallic Corrosion, NACE, Vol. 312 (1974).
5. Pickering, H. M. and Frankenthal, R. P. Proceedings of the Fifth International Congress on Metallic Corrosion, NACE, Vol. 261 (1974).
6. Sato, N. Passivity and its Breakdown on Iron and Iron Base Alloys, USA-Japan Seminar, R. W. Staehle and H. Okada, NACE, Vol. 1 (1976).
7. Mattsson, E. Proceedings of the 6th European Congress on Metallic Corrosion, The Society of Chemical Industry, London, Vol. 219 (1977).
8. Olefjord, I. in S. Nilsson (Ed.), Proc. 6th Scand. Corrosion Congress, Gothenburg, 1971, Swedish Corrosion Institute, Stockholm, p. 11-1 (1971).
9. Fischmeister, H. and Olefjord, 1. Monatsh. Chem., Vol. 102, p. 1486 (1971).
10. Olefjord, I. Scand. J. Metall., Vol. 3, p. 129 (1974).
11. Olefjord, I. Corrosion Science, Vol. 15, p. 687 (1975).
12. Olefjord, I. and Fischmeister, H. Corrosion Science, Vol.
15, p. 697 (1975).
13. Olefjord, I. and Elfstrom, B-O. Proc. 8th Int. Symp. on Reac- tivity of Solids, Plenum, New York, p. 791 (1977).
14. Olefjord, I. and Elfstrom, B-O. in T. P. Hoar (Ed.), Proc. 6th Eur. Congress on Metallic Corrosion, London, 1977, Soci- ety of Chemical Industry, London, p. 21 (1977).
15. Elfstrom, B-O and Olefjord, I. Physica Scripta, Vol. 16, p. 436 (1977).
16. Olefjord, 1. in Larinkari, J. (Ed.), Proc. 8th Scand. Corro- sion Congress, Helsinki University of Technology, Labora- tory of Corrosion Science and Technology, Helsinki, p. 349 (1978).
17. Elfstrom, B-O. and Olefjord, I. in J. Larinkari (Ed.), Proc. 8th Scand. Corrosion Congress, Helsinki, 1978, Helsinki University of Technology, Laboratory of Corrosion Science and Technology, Helsinki, p. 312 (1978).
18. Olefjord, I. Materials Science and Engineering, Vol. 42, p. 161 (1980).
19. Elfstrom, B.O. Materials Science and Engineering, Vol. 42, p. 173 (1980).
20. Sugimoto, K. and Sawada, Y. Corrosion Science, Vol. 17, p. 425 (1977).
21. Castte, J. E. and Clayton, C. R. Corrosion Science, Vol. 17, p. 7 (1977).
22. Yani, A. E., Lumsden, J. B., and Staehle, R. W. J. Electro- chem. Soc., Vol. 124, p. 490 (1977).
23. Leygraf, C., Hultquist, G., Olefjord, I., Elfstrom, B-O., Knyasheva, V. M., Plaskeyev, A. V. and Kolotyrkin, Ya. M.
Corrosion Science, Vol. 19, p. 343 (1979).
24. Okamoto, G. Corrosion Science, Vol. 13, p. 471 (1973).
25. Barnes, G. J., Aldag, A. W., and Jerner, R. C. J. Electro- chem. Soc.: Electrochemical Science and Technology, Vol.
119, p. 684 (1972).
26. Charbounier, J. C., Maitrepierre, Ph., Noual, P., and Namdar-Irani, R. Proc. of the 7th Intern. VAC Congress and the 3rd Intern. Conf. Solid Surfaces, Vienna (1977).
27. Hashimoto, K., Asami, K., and Teramoto, K. Corrosion Science, Vol. 19, p. 3 (1979).
28. Ogawa, H., Omata, H., Itok, I., and Okada, H. Corrosion, NACE (Houston), Vol. 34, p. 52 (1978).
29. Olefjord, I., Leijon, W., and Jelvestam, U. Applications of Surface Science, Vol. 6, p. 241 (1980).
30. Olefjord, I. J. Electron Spectroscopy and Related Phe- nomena, Vol. 5, p. 401 (1974).
31. Marcus, P., Oudar, J., and Olefjord, I. To be published in Surface and Interface Analysis.
32. Marcus, P., Oudar, J., and Olefjord, 1. J. De Microscopie et de Spectroscopie electriques, Vol. 4, p. 63 (1979).
33. Fadley, C. S. in G. Somorjai and J. McCaldin, Progress in Solid State Chemistry, Pergamon Press, New York, Vol. 11, p. 265 (1976).
34. Powell, C. J. Surface Science, Vol. 44, p. 29 (1974).
35. Nilsson, P. and Hakansson, B. in J. Larinkari (ed.) Proc. 8th Scand. Corrosion Congress, Helsinki, 1978, Helsinki University of Technology, Laboratory of Corrosion Science and Technology, Helsinki, p. 41 (1978).
36. Elfstrom, B-O., and Olefjord, 1. To be published.
37. Galvele, J. R., Lumsden, J. B., and Staehle, R. W. J. Electro- chemical Soc.: Electrochemical Science and Technology, Vol. 125, p. 1204 (1978).
38. Sugimoto, K. and Matsuda, S. Materials Science and Engi- neering, Vol. 42, p. 181 (1980).
39. Shubat, G. J. Metals Handbook, ASM, Ohio, Vol. 8, p. 421 and 426 (1973).
40. Hakkarainen, T. and Olefjord, I. To be published.
