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Effects of alloying elements, Cr, Mo and N on repassivation characteristics of stainless steels using the abrading electrode technique
Materials Chemistry and Physics 99 (2006) 224–234
Effects of alloying elements, Cr, Mo and N on repassivation characteristics of stainless steels using the abrading electrode technique Jae-Bong Lee ∗ School of Advanced Materials Engineering, Kookmin University 861-1, Jeongneung-dong, Sungbuk-gu, Seoul 136-702, Korea Received 16 April 2005; received in revised form 8 August 2005; accepted 21 October 2005
Abstract Cr, Mo and N are important alloying elements for passive film formation and repassivation but their roles in passive film still need explainable. In order to investigate effects of alloying elements on the stability of passive film and its repassivation, ferritic stainless steels such as Fe–Cr and Fe–Cr–Mo alloys and austenitic stainless steels such as type 316L, and type 316LN were examined. The electrochemical characteristics of the passive film were investigated by in situ dc and ac electrochemical methods. The resistance to localized corrosion is believed to have much to do with the repassivation of the passive film and its stability. The effects of alloying elements on the current transients were systematically examined by using the abrading electrode technique and electrochemical impedance spectroscopy (EIS). The experimental results were analyzed in order to elucidate the relationship between passive film stability, repassivation, and alloying elements. © 2005 Elsevier B.V. All rights reserved. Keywords: Repassivation; Passive film; Abrading electrode technique; Electrochemical impedance spectroscopy (EIS); Charge transfer resistance (Rct )
1. Introduction Stainless steels have excellent corrosion resistance, resulting from thin and protective passive films, which prevent metals from reacting with corrosive environments. The influence of alloying elements on corrosion resistance is related with the stability of the passive film and its repassivation kinetics. The breakdown of this film can occur by chemical or by mechanical attack in corrosive media resulting in the localized corrosion. Pits formed by film breakdown can be the potential initiation sites for various localized corrosion such as pitting, crevice corrosion and stress–corrosion cracking [1,2]. Thus, in order to prevent localized corrosion, it is important for stainless steels to have a stable passive film with rapid repassivation even in severe corrosive environments. It is believed that the stability of passive films and their repassivation kinetics are subject to the influence of metallurgy, applied repassivation potentials, pH and chloride ion concentration in aqueous solutions [3–9]. The present study is designed to explain the relationship between the passive film stability, repassivation kinetics and alloying elements, Cr, Mo and N. In order to understand the effect of alloying elements on the corrosion resistance to corrosion of various stainless steels, ∗
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in situ dc and ac electrochemical techniques were carried out. The influence of alloying elements on repassivation was systematically investigated using the abrading electrode technique [8,9] and ac impedance spectroscopy. Furthermore, the effects of the chloride ion concentration and the applied potential were also investigated. 2. Experimental details 2.1. Materials and preparations Fe–xCr (x = 0, 13, 18, 25, 30, 40 and 100 wt.%) alloys and Fe–18Cr–xMo (x = 1, 4, and 6 wt.%) alloys were prepared by melting in a vacuum arc melting furnace back-filled with argon, and subsequently remelting numerous times to ensure proper mixing. Austenitic stainless steels such as type 316L and type 316LN with the variation of N concentrations from 0 to 0.15 wt.% were prepared in the form of ingots by melting in vacuum induction melting furnace. After all samples were solution-treated at 1050–1270 ◦ C to avoid segregation, they were hot-rolled by 50–83% reduction in thickness, annealed at 1050–1100 ◦ C, followed by water quenching. Each specimen was wet-ground up to 2000 grit SiC paper followed by fine polishing with a suspension of 0.05 m particles of alumina. The samples were degreased with ethyl alcohol and acetone
2.2. Electrochemical measurements The 0.1N H2 SO4 + xNaCl (x = 0, 0.1 and 0.6N) acidic solution (pH 1.6) were prepared with distilled water and reagent grade H2 SO4 and NaCl. The solution was de-aerated with high purity Ar before testing and kept under Ar atmosphere
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the pitting potential, decreasing the passive current density and the critical anodic current density. Fig. 2 shows potentiodynamic polarization curves for pure Fe, pure Cr, and Fe–xCr (x = 13, 18, 25, 30, and 40 wt.%) alloys in aqueous 0.1N H2 SO4 without and with Cl− ion. Fig. 2(a) shows results for anodic polarization curves for Fe–Cr alloys with varying Cr content in aqueous 0.1N H2 SO4 without any Cl− ion, representing two kinds of passive potential regions which consisted of the primary low current density plateau and the secondary high current density plateau [13]. As Cr content increased, the primary passive current density decreased, exhibiting the secondary passive region disappeared. However, in the case of pure Fe, the high current density region was
Fig. 1. Apparatus for abrading electrode used in the study of repassivation kinetics of free electrode surface. during testing. A platinum wire and a saturated calomel electrode (SCE) were used as the counter and reference electrodes, respectively. Potentiodynamic polarization experiments were performed on specimens using a CMS 105B potentiostat controlled by a computer software (CMS 105). An abrading electrode technique was employed to obtain information on the repassivation kinetics of oxide films repassivated on the surface of specimens after exposure of the bare surface to the solution. The electrode was abraded with an SiC disc while immersed in the electrolyte to acquire current transients as shown in Fig. 1. The specimen was mounted in the center of a polyethylene resin (diameter = 1 cm) with a exposed area of 8 × 10−2 cm2 which was welded to a current collector of copper wire. The specimen surface was renewed by abrading with the SiC disc sintered from 1200 grit. The disc was fixed to the rotating shaft operated by dc motor. A saturated calomel electrode (SCE) was used as the reference electrode, located nearest to the working electrode. Finally a graphite was used as the counter electrode. During the abrading, the specimen was instantaneously raised by a mechanical spring. The resulting current transient was acquired at a constant applied potential from the moment just after interrupting the abrading action on the specimen. In order to evaluate the electrochemical behavior of repassivated films, ac impedance spectroscopy measurements were performed by using a potentiostat and a lock-in amplifier (SR810) with an electrochemical impedance software (CMS300) under static conditions upon completion of the abrading tests. Measurements were taken by superimposing an ac voltage 10 mV amplitude on the measured rest potential of the abraded specimen with the frequency range from 0.01 Hz to 10 kHz.
3. Results and discussion 3.1. Effect of Cr content on the passivity and the repassivation Stainless steels basically derive their passive characteristics from alloying with Cr. Alloying element, Cr is known to increase resistance to pitting by the increase of the pitting potential in the noble direction, reducing the passive current density. This is illustrated clearly by the data for a series of Fe–Cr–Ni alloys in a sulfuric acid solution [10], and for Fe–Cr alloys in a chloride solution [11], respectively. A.J. Sedriks summarized the effect of alloying elements in stainless steels on the anodic polarization curve [12]. According to his schematic summary, Cr increases
Fig. 2. Potentiodynamic polarization curves for Fe–Cr alloys in deaerated 0.1N H2 SO4 without and with Cl− ion according to the variation of Cr content: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.1N NaCl.
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Fig. 3. Decay of current density with time during repassivation of Fe–Cr alloys at the applied potential of −0.25 V in deaerated 0.1N H2 SO4 .
observed up to +900 mV above which a steep decrease in current followed by a low current density region was noted. Since this region corresponds to the secondary passive potential region of the anodic curves of Fe–Cr alloys, in addition to Cr oxide, the formation of Fe oxide seems to play some role on the secondary passive film of Fe–Cr alloys. However, since the primary passive potential region of Fe–Cr alloys agrees with the passive potential region of pure Cr and furthermore, with increasing Cr content, the polarization behaviors of Fe–Cr alloys approximates that of pure Cr, the Cr oxide on the surface of specimen plays the most important role on the formation of primary passive film. In order to investigate the effect of Cl− ion on the passive film, potentiodynamic curves for Fe–Cr alloys in aqueous 0.1N H2 SO4 with 0.1N NaCl is shown in Fig. 2(b). Pure Cr showed the similar kinds of polarization behaviors irrespective of Cl− ion concentrations while pure Fe did not display the secondary passive region due to addition of Cl− ion, representing Cr oxide film has the more resistance to corrosion than Fe oxide film in the chloride ion containing solutions. Therefore, based on the results of polarization curves, the passivity of Fe–Cr alloys may be mainly attributed to Cr oxide film whereas Fe oxide film has the less influence on the passivation behavior in the chloride ion containing acidic solutions The repassivation kinetics of Fe–Cr alloys can be expressed as follows. i = At −n
(1)
where i denotes the anodic current density, A is the constant, t is the time, and n is the repassivation rate parameter representing the slope of log i − log t plot. Based on Eq. (1), the value of n can be considered to be an indirect measure of the growth of the oxide on the bare surface as a function of applied anodic potential. Fig. 3 shows the potentiostatic current density decays plotted in log i versus log t obtained from the moment just after interrupting the abrading for Fe–Cr alloys applied at −0.25 V
in deaerated 0.1N H2 SO4 . Repassivation current density, i with time, t was observed to obey the linear relationship of log i versus log t [14–16]. The anodic current density of pure Fe was nearly constant maintaining high values with time on a logarithmic scale while that of pure Cr showed the lowest among pure Fe, pure Cr and Fe–Cr alloys. As the Cr content increased in Fe–Cr alloys, repassivation current densities decreased, slope −n of log i = k − n log t approaching to that of pure Cr. Pure Fe did not show any repassivation behavior at −0.25 V, in deaerated 0.1N H2 SO4 . These demonstrate that alloying element Cr in Fe–Cr alloys plays an important role in the rapid repassivation after fracturing the passive film Fig. 4 shows the decay of the current density with time in logarithmic scale during the repassivation of pure Cr with varying applied potential and Cl− ion concentration in deaerated 0.1N H2 SO4 solution. In the case of 0.1N H2 SO4 solution without Cl− ion, regardless of applied potential, pure Cr samples were completely repassivated with increasing time, showing that repassivation current densities increased with increasing applied potential in the passive potential range (Fig. 4(a)). However, in 0.1N H2 SO4 solution containing 0.1N Cl− ion, repassivation did not maintain at applied potential of 0.6 V, showing n = 0 after t = 200 s (Fig. 4(b)). As shown in Fig. 4(c), when the chloride ion concentration was increased to 0.6N NaCl and the anodic potential of 0.3 V was applied, an ascending anodic current density was observed, after t = 20 s, indicating the growth of stable pits [17]. Thus, in the case of absence of chloride ion, pure Cr showed the same repassivation kinetics irrespective of applied potential but in the presence of the chloride ion, even pure Cr was not able to become completely repassivated, as chloride ion concentration and applied potentials increased. In order to investigate effects of Cr content and applied potential on the corrosion resistance, ac impedance spectra for pure Cr and Fe–Cr alloys were measured at the open circuit potential in deaerated 0.1 M Na2 SO4 solution just after repassivation during 200 s. at various applied potentials −0.25, 0.3 and 0.6 V in deaerated 0.1N H2 SO4 solution. The passive oxide films formed by repassivation were analyzed by using the equivalent circuit model [18] (Fig. 5), which admits a one time constant Warburg diffusion model and the non-linear least squares (NLLS) fitting method. The charge transfer resistance (Rct ) of the repassivated film tended to increase with increasing Cr content and applied potential (Fig. 6(a)) while the capacitance of passive film (Cox ) tended to decrease with the increase in Cr content (Fig. 6(b)). The present results on the improvement of repassivated film with increasing Cr content may be explained by the investigation on the role of Cr on the passive film of Fe–Cr alloys in sulfuric acid carried out by J.A.L. Dobbelaar et al. [19]. They found that the porosity in a passive film decreases with increasing Cr content and/or applied potential, improving the protective quality of the film. 3.2. Effect of Mo content on the passivity and the repassivation Fig. 7 shows potentiodynamic polarization curves for Fe–18Cr–xMo (x = 1, 4, and 6 wt.%) alloys in aqueous 0.1N
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Fig. 4. Decay of current density with time for during repassivation of pure Cr at varying applied potential and Cl− concentration in deaerated 0.1N H2 SO4 solution: (a) 0.1N H2 SO4 , (b) 0.1N H2 SO4 + 0.1N NaCl and (c) 0.1N H2 SO4 + 0.6N NaCl.
H2 SO4 without and with Cl− ion concentrations of 0.1 N. In aqueous 0.1N H2 SO4 without Cl− ion, the increase in Mo content showed approximately same anodic polarization curves, compared with that of Fe–18Cr alloy without any Mo (Fig. 7(a)). However, in aqueous 0.1N H2 SO4 with Cl− ion, as Mo content increased, the pitting potential was raised, representing the passive potential region was extended in width (Fig. 7(b)). At the same time, the critical current density for passivation, ic decreased with increasing Mo content [20]. This suggests that the higher Mo content promotes the easier passivation and the more stable passivating oxide film. According to the results of Auger Electron Spectroscopy (AES) by Ogawa et al. [21] and X-ray Photo-electron Spectroscopy (XPS) by Hashimoto et al. [22], Mo can be interpreted
to play a significant role not only to promote the increment of the Cr content in the passive film but also to lead to the formation of stable and homogeneous passive films due to the decrease in the activity of surface sites. Montemor et al. [23] made analytical experiments for the passive film shown as a duplex structure for the oxides in stainless steels. Based on their results for AES and XPS, Mo was observed to be present in the internal Cr rich layers as well as at the Cr oxide/metal interface, resulting in an enrichment of Cr with the assistance of Mo in the internal region of the passive film. They also suggested that Mo introduces changes in the properties and the defect structure of the oxide film. Those experimental results associated with both the increase in pitting potentials and the extension of passive potential region in width were observed in Fig. 7, as Mo content increased.
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slope is in good agreement with minute pits on the metal surface observed after abrading Fe–18Cr alloy at −0.25 V in deaerated 0.1N H2 SO4 + 0.6N NaCl solution. In addition, Fig. 9(a) and (b) shows that the increase in Mo content causes the faster repassivation by reducing peak current density transients regardless of the variation of repassivation rate. Fig. 10 shows impedance measurement results in Nyquist plots for Fe–18Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys carried out just after performing abrading electrode test at an applied potential of −0.25 V. The ac impedance presented here were measured at the open circuit potential, Eoc in 0.1N H2 SO4 and 0.1N H2 SO4 + 0.6N NaCl, respectively. The diameter of the semicircle i.e. the charge transfer resistance, Rct increases with increasing Mo content. The increase in Rct with Mo content may be interpreted in the fact that the presence of Mo in repassivated film may shift the alloy from the active region to the pre-passive region, and form a pre-cursor film for passivation, since Rct of Fe–18Cr alloy without Mo is quite low. J.H.
Fig. 5. Equivalent circuit model for simulation of Fe–Cr alloys [18].
Fig. 8 shows the current density transients of Fe–18Cr and Fe–18Cr–xMo (x = 1, 4, 6 wt.%) alloys with time carried out just after performing abrading electrode test in deaerated 0.1N H2 SO4 solution at an applied potential of −0.25 V. Current density transient peak values decreased as Mo content increased, representing that addition of Mo leads to a significant decrease in the transient heights. Fig. 9 shows the potentiostatic current density decays plotted in log i versus log t obtained from the moment just after interrupting the abrasion of Fe–18Cr–xMo alloys applied at −0.25 V in deaerated 0.1N H2 SO4 (Fig. 9(a)) and 0.1N H2 SO4 + 0.6N NaCl (Fig. 9(b)). Repassivation current density with time was observed to obey the linear relationship of log i versus log t. In deaerated 0.1N H2 SO4 both Fe–18Cr and Fe–18Cr–xMo (x = 1,4, and 6 wt.%) alloys showed all anodic current densities continuously decreased with time in logarithmic scale. However in deaerated 0.1N H2 SO4 + 0.6N NaCl, anodic current density of Fe–18Cr alloy descended linearly with time in logarithmic scale and then the value of current density approached to nearly constant while those of Fe-18Cr–xMo alloys continuously descended linearly with time in logarithmic scale. The fall of anodic current density can be explained by the fact that the repassivation rate of passivating oxide film dominated over its dissolution rate [24]. The slope, −n of log i versus log t close to −1 is typical for passive films growing without the superimposed dissolution of passive films while 0 or positive slopes represent uniform corrosion or pitting corrosion, respectively. Therefore, the variation of a slope of Fe–18Cr with time, at −0.25 V in deaerated 0.1N H2 SO4 + 0.6N NaCl, imply the possibility of the dissolution of repassivated layers of Fe–18Cr after around 100 s whereas slopes of Fe–18Cr–xMo alloys close to −1 with time indicate the formation of the sound and protective film without the dissolution of passive layers during abrading tests. This possibility of pitting corrosion due to the variation of
Fig. 6. Changes of Rct and Cox of Fe–Cr alloys obtained from the impedance spectra with varying applied potentials and Cr content: (a) Rct and (b) Cox .
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Fig. 8. Current density transients curve of Fe–18Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys in deaerated 0.1N H2 SO4 solution at an applied potential of −0.25 V.
and stabilization of Cr oxide film [16]. It has also been suggested that Mo retards the corrosion process by adsorption [30], by formation of Mo compounds [31], by synergistic interaction of Mo ions with other oxides of the passive film [20,32] and by elimination of the active surface sites through formation of Mo oxides [22]. Even though there are analytical results showing the presence of Mo, in other studies, Mo was not found in the passive film [22,27,28]. Therefore, the satisfactory explanation for the role of Mo on corrosion resistance is still a subject of discussion. Based on these present experimental results, the increase in Mo content raises the pitting potentials, producing a more stable passive film and increasing the charge transfer resistance. Furthermore, the addition of Mo facilitates the faster repassivation by reducing peak current density transients,
Fig. 7. Potentiodynamic polarization curves for Fe–18Cr–xMo alloys in deaerated 0.1N H2 SO4 without and with Cl− ion according to the variation of Mo content: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.1N NaCl.
Gerretsen and J.H.W. de Wit investigated the role of Mo on the passivation of Fe–Cr alloys [25]. According to their results, Mo increases the dissolution rate of the alloys, but the increase is compensated by the faster formation of the passive film, thus increase the resistance to corrosion of the Fe–Cr alloys, which is in agreement with the present experimental results showing that the addition of Mo to Fe–18Cr causes the faster repassivation during abrading and the higher charge transfer resistance measured by ac impedance spectroscopy after abrading test. The beneficial effect of Mo in the corrosion resistance of stainless steels has been known to be attributed to the following several factors. These include an enrichment of Mo in the passive film [20,26] or just below the passive film [27,28], enrichment of Cr in the oxide layer [28,29], thickening of the passive film
3.3. Effect of N content on the passivity and the repassivation Fig. 11 shows potentiodynamic polarization curves for type 316L (0 wt.% N), type 316LN (0.042 wt.% N), type 316LN (0.1 wt.% N) and type 316LN (0.15 wt.% N) austenitic stainless steels in deaerated 0.1N H2 SO4 with varying Cl− ion concentration. As shown in Fig. 11(a), N addition to type 316LN in deaerated 0.1N H2 SO4 without Cl− ion did not make any noticeable difference except for somewhat reducing the critical current density for passivation, ic . Even though Cl− ion concentration increases up to 0.6N, the addition of N seems to have the anodic polarization curves of type 316LN stainless steels unchanged. However, the anodic polarization curve of type 316L without N in deaerated 0.1N H2 SO4 + 0.6N NaCl solution, shows the much higher passive current density, ipass and critical anodic current density for passivation, ic than type 316LN with N, indicating that N addition promotes the easier passivation and the more stable passivating oxide film. In the similar manner of Mo, the
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linearly with time in logarithmic scale but once the current densities reached minimum current values, the current densities were observed to ascend. This ascending anodic current density can be attributed to the formation and growth of pits due to the passive film breakdown, depending upon applied anodic potential and chloride ion concentration [17]. However the anodic current densities of type 316LN (0.042 wt.% N) and type 316LN (0.15 wt.% N) continuously descended linearly with time in logarithmic scale irrespective of applied potentials, indicating that the repassivation rate of passivating oxide film dominated over its dissolution rate on the bare surface. The role of N in corrosion resistance of stainless steels can be explained by suppressing the dissolution of oxide film and by facilitating the repassivation of the passive film, which cause a decrease in the height of current transients.
Fig. 9. Decay of anodic current density with time for repassivation of Fe–Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys at −0.25 V in: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.6N NaCl.
addition of N leads to increase the pitting potentials, widen the passive potential region, and reduce the critical anodic current density, promoting the faster formation of passive films and a more stable passive film. These results seems to be related with those reported by Kamachi Mudali et al. [33] that pit initiation was delayed and pit growth was significantly decreased with the addition of N to pure iron. Fig. 12 shows the potentiostatic current density decays plotted in log i versus log t for type 316L (0 wt.% N), type 316LN (0.042 wt.% N) and type 316LN (0.15 wt.% N) austenitic stainless steels obtained at applied potentials of 0, 0.4, and 0.6 V in deaerated 0.1N H2 SO4 + 0.1N NaCl. Repassivation current densities of type 316L without N at 0.4 and 0.6 V initially descended
Fig. 10. Impedance spectra in Nyquist plot for Fe–18Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys carried out at Eoc just after performing abrading electrode test at an applied potential of −0.25 V in: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.6N NaCl.
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Fig. 11. Potentiodynamic polarization curves for 316L and 316LN with different N content in deaerated 0.1N H2 SO4 with varying Cl− ion concentration: (a) 0.1N H2 SO4 , (b) 0.1N H2 SO4 + 0.1 NaCl and (c) 0.1N H2 SO4 + 0.6N NaCl.
Fig. 13 shows impedance spectra in Bode plot for type 316L (0 wt.% N), type 316LN (0.042 wt.% N), and type 316LN (0.15 wt.% N) carried out at the open circuit potential, Eoc after performing potentiostatic current density transient measurements at applied potentials of 0 and 0.6 V in deaerated 0.1N H2 SO4 . The charge transfer resistance, Rct of type 316LN increased with N addition by 10 times compared with that of type 316L without N. Fig. 13(b) shows charge transfer resistance of repassivated film increased with the increase in applied potential from 0 to 0.6 V in deaerated 0.1N H2 SO4 . It is suggested that the passive film formed at higher applied potential is more protective than that produced at lower applied potential within the
passive potential region of deaerated 0.1N H2 SO4 solution. This indicates that the higher potentials apply in the passive potential region, the more stable oxide films form in deaerated sulfuric acid solution. The influence of N on the corrosion resistance of stainless steel has received much attention in recent years. There is a general consensus that N additions improve the resistance to pitting corrosion, increasing the pitting potential in aqueous chloride solutions. The present experimental work shows that N also accelerate repassivation. In terms of the role of N in the resistance to corrosion, Osozawa and Okada [34] attributed the improved pitting
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Fig. 12. Decay of anodic current densities with time for repassivation kinetics of 316L, 316LN (0.042 wt.% N) and 316LN (0.15 wt.% N) with various applied potentials in deaerated 0.1N H2 SO4 + 0.1N NaCl solution: (a) 316L; (b) 316LN (0.042 wt.% N) and (c) 316LN (0.15 wt.% N).
resistance of N-containing stainless steels to the formation of ammonium ions at pit sites, which increased the pH. Newman and Shahrabi [35] classified the effects of N on pitting corrosion resistance into three steps, namely (1) consumption of acid in pit nuclei by nitrogen dissolution, i.e. N + 4H+ + 3e → NH4 + , (2) enrichment of N on the passivated surface, and (3) enrichment of N on active surfaces. Newman et al. [36] also observed enrichment of Mo and N at the passive film/metal interface and attributed this enrichment to the prevention of the dissolution of the substrate following destruction of the passive film. JargeliusPettersson [37] found the N additions increased the critical temperature for pitting corrosion, indicating a synergistic interaction between Mo and Olefjord and Wegrelius [38] observed
that the high N-containing steel was repassivated while the pits were growing on the low N-containing steel, suggesting a model for a synergistic effect between Mo and N. However, only a few studies concerning the effect of Mo and N on the repassivation kinetics has been reported, and it is proposed that N-containing stainless steels without Mo, did not show any influence of the N content on the polarization behavior [39,40]. Therefore, even though based upon the present studies, alloying elements, Mo and N promote the passivation by formation of more stable passive film and the repassivation by increasing the formation rate of passive film, the exact mechanisms of the role of Mo and N on the repassivation kinetics and the exact interaction between N and Mo remains still to be clarified.
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with increasing applied potential within the passive potential region. 2. The increase in Mo content raises the pitting potentials and reduces the critical anodic current density for passivation in aqueous 0.1N H2 SO4 with Cl− ion. The addition of Mo facilitates passivation, produces the more stable passive film, causes the faster repassivation by reducing peak current density transients and finally promotes the resistance to corrosion by increasing the charge transfer resistance. 3. Compared type 316L and type 316LN stainless steels, type 316LN to which N is added shows a much lower passive current density and critical anodic current density for passivation. In the similar manner of the addition of Mo, the addition of N increases the repassivation rate of the oxide film, reducing the corrosion rate. Acknowledgment The author gratefully acknowledges the financial support by Korea Science and Engineering Foundation (KOSEF). References [1] [2] [3] [4] [5]
[6]
Fig. 13. Impedance spectra in Bode plot for 316L, 316LN (0.042 wt.% N) and 316LN (0.15 wt.% N) carried out at Eoc , after performing potentiostatic current density transient measurements in deaerated 0.1N H2 SO4 at applied potentials of (a) 0 V and (b) 0.6 V.
[7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
4. Conclusions 1. The corrosion potentials and the passive current densities of Fe–Cr alloys were decreased with the increase in Cr content. The passivity of Fe–Cr stainless steels was mainly attributed to primary Cr oxide film. In addition to Cr oxide film, Fe oxide may have something to do with secondary passive film. Addition of Cl− ion to 0.1N H2 SO4 solution resulted in the contracted passive potential region and the increased passive current density. The increase in Cr content of Fe–Cr alloys increased both the repassivation rate and charge transfer resistance, Rct . In the case of addition of Cl− ion, the tendency of suppressing repassivation rate increased
[20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]
T.R. Beck, Corrosion 30 (1974) 408. J.C. Scully, Corros. Sci. 15 (1975) 207. R.C. Newman, Corros. Sci. 25 (1985) 331. P.I. Marshall, G.T. Burstein, Corros. Sci. 23 (1983) 1219. F. Falkenberg, I. Olefjord, in: M. Seo, E. MacDougall, H. Takahashi, R.G. Kelly (Eds.), Passivity and Localized Corrosion, The Electrochemical Society, Inc., Pennington, NJ, 1999, p. 404. Y. Virtanen, S. Olefjord, H. Bohni, in: M. Seo, E. MacDougall, H. Takahashi, R.G. Kelly (Eds.), Passivity and Localized Corrosion, The Electrochemical Society, Inc., Pennington, NJ, 1999, p. 533. J.S. Kim, P.J. Xiang, K.Y. Kim, Corrosion 61 (2005) 174. J.-D. Kim, S.-I. Pyun, Corros. Sci. 38 (1996) 1093. S.-I. Pyun, M.-H. Hong, Electrochim. Acta 37 (1992) 2437. K. Osozawa, H.J. Engell, Corros. Sci. 6 (1966) 389. J. Horvath, H.H. Uhlig, J. Electrochem. Soc. 115 (1968) 791. A.J. Sedriks, Corrosion 42 (1986) 376. A. Atrens, B. Barou, M. Mantel, J. Electrochem. Soc. 144 (1997) 3697. T.R. Beck, R.C. Alkire, J. Electrochem. Soc. 131 (1984) 1662. P.I. Marshall, G.T. Burstein, Corros. Sci. 23 (1983) 1219. J.W. Schultze, M.M. Lohrengel, D. Ross, Electrochim. Acta 28 (1993) 973. T.H. Nguyen, R.T. Foley, J. Electrochem. Soc. 126 (1979) 1855. CMS300 Electrochemical Impedance Spectroscopy System Operator’s Manual, Gamry Instrument Inc., PA, USA, 1994, pp. 7–22. J.A.L. Dobbelaar, E.C.M. Herman, J.H.W. de Wit, Corros. Sci. 33 (1992) 765. K. Sugimoto, Y. Sawada, Corros. Sci. 17 (1977) 425. H. Ogawa, H. Omata, I. Itoh, H. Okada, Corrosion 34 (1978) 53. K. Hashimoto, K. Asami, K. Teramoto, Corros. Sci. 19 (1979) 3. M.F. Montemor, A.M.P. Simoes, M.G.S. Ferreira, M. Da Cunha Belo, Corros. Sci. 41 (1999) 17. J.-J. Park, S.-I. Pyun, W.-J. Lee, H.-P. Kim, Corrosion 55 (1999) 380. J.H. Gerretsen, J.H.W. de Wit, Electrochim. Acta 36 (1991) 1465. C. Leygraf, G. Huktquist, I. Olefjord, B.O. Elfstrom, V.M. Knayazheva, A.V. Paskeyer, Y.M. Kolotyrkin, Corros. Sci. 19 (1979) 343. I. Olefjord, B.O. Elfstrom, Corrosion 38 (1982) 46. M.W. Tan, E. Akiyama, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci. 37 (1995) 1289. H. Habazaki, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci. 33 (1992) 225. K. Sugimoto, Y. Sawada, Corrosion 32 (1976) 347.
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[31] C.R. Clayton, J.C. Lu, Corros. Sci. 27 (1989) 881. [32] E. Akiyama, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci. 38 (1996) 1281. [33] U. Kamachi Mudali, B. Reynders, M. Stratmann, Corros. Sci. 41 (1999) 179. [34] K. Osozawa, N. Okada, Passivity and its Breakdown on Iron and Iron-Base Alloys, NACE, Houston, USA, 1976, p. 135. [35] R.C. Newman, T. Shahrabi, Corros. Sci. 27 (1987) 827.
[36] R.C. Newman, Y.C. Lu, R. Bandy, C.R. Clayton, Proceedings of the Ninth International Congress on Metallic Corrosion, vol. 1, National Research Council, Canada, Toronto, 1984, p. 394. [37] R.F.A. Jargelius-Pettersson, Corros. Sci. 41 (1999) 1639. [38] I. Olefjord, L. Wegrelius, Corros. Sci. 38 (1996) 1203. [39] M. Janik-Czachor, E. Lunarska, Z. Szklaraska-Smialowska, Corrosion 31 (1975) 349. [40] A. Sadough Vanini, R. Chikhi, P. Marcus, Corros. Sci. 36 (1994) 1825.
Effects of alloying elements, Cr, Mo and N on repassivation characteristics of stainless steels using the abrading electrode technique Jae-Bong Lee ∗ School of Advanced Materials Engineering, Kookmin University 861-1, Jeongneung-dong, Sungbuk-gu, Seoul 136-702, Korea Received 16 April 2005; received in revised form 8 August 2005; accepted 21 October 2005
Abstract Cr, Mo and N are important alloying elements for passive film formation and repassivation but their roles in passive film still need explainable. In order to investigate effects of alloying elements on the stability of passive film and its repassivation, ferritic stainless steels such as Fe–Cr and Fe–Cr–Mo alloys and austenitic stainless steels such as type 316L, and type 316LN were examined. The electrochemical characteristics of the passive film were investigated by in situ dc and ac electrochemical methods. The resistance to localized corrosion is believed to have much to do with the repassivation of the passive film and its stability. The effects of alloying elements on the current transients were systematically examined by using the abrading electrode technique and electrochemical impedance spectroscopy (EIS). The experimental results were analyzed in order to elucidate the relationship between passive film stability, repassivation, and alloying elements. © 2005 Elsevier B.V. All rights reserved. Keywords: Repassivation; Passive film; Abrading electrode technique; Electrochemical impedance spectroscopy (EIS); Charge transfer resistance (Rct )
1. Introduction Stainless steels have excellent corrosion resistance, resulting from thin and protective passive films, which prevent metals from reacting with corrosive environments. The influence of alloying elements on corrosion resistance is related with the stability of the passive film and its repassivation kinetics. The breakdown of this film can occur by chemical or by mechanical attack in corrosive media resulting in the localized corrosion. Pits formed by film breakdown can be the potential initiation sites for various localized corrosion such as pitting, crevice corrosion and stress–corrosion cracking [1,2]. Thus, in order to prevent localized corrosion, it is important for stainless steels to have a stable passive film with rapid repassivation even in severe corrosive environments. It is believed that the stability of passive films and their repassivation kinetics are subject to the influence of metallurgy, applied repassivation potentials, pH and chloride ion concentration in aqueous solutions [3–9]. The present study is designed to explain the relationship between the passive film stability, repassivation kinetics and alloying elements, Cr, Mo and N. In order to understand the effect of alloying elements on the corrosion resistance to corrosion of various stainless steels, ∗
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in situ dc and ac electrochemical techniques were carried out. The influence of alloying elements on repassivation was systematically investigated using the abrading electrode technique [8,9] and ac impedance spectroscopy. Furthermore, the effects of the chloride ion concentration and the applied potential were also investigated. 2. Experimental details 2.1. Materials and preparations Fe–xCr (x = 0, 13, 18, 25, 30, 40 and 100 wt.%) alloys and Fe–18Cr–xMo (x = 1, 4, and 6 wt.%) alloys were prepared by melting in a vacuum arc melting furnace back-filled with argon, and subsequently remelting numerous times to ensure proper mixing. Austenitic stainless steels such as type 316L and type 316LN with the variation of N concentrations from 0 to 0.15 wt.% were prepared in the form of ingots by melting in vacuum induction melting furnace. After all samples were solution-treated at 1050–1270 ◦ C to avoid segregation, they were hot-rolled by 50–83% reduction in thickness, annealed at 1050–1100 ◦ C, followed by water quenching. Each specimen was wet-ground up to 2000 grit SiC paper followed by fine polishing with a suspension of 0.05 m particles of alumina. The samples were degreased with ethyl alcohol and acetone
2.2. Electrochemical measurements The 0.1N H2 SO4 + xNaCl (x = 0, 0.1 and 0.6N) acidic solution (pH 1.6) were prepared with distilled water and reagent grade H2 SO4 and NaCl. The solution was de-aerated with high purity Ar before testing and kept under Ar atmosphere
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the pitting potential, decreasing the passive current density and the critical anodic current density. Fig. 2 shows potentiodynamic polarization curves for pure Fe, pure Cr, and Fe–xCr (x = 13, 18, 25, 30, and 40 wt.%) alloys in aqueous 0.1N H2 SO4 without and with Cl− ion. Fig. 2(a) shows results for anodic polarization curves for Fe–Cr alloys with varying Cr content in aqueous 0.1N H2 SO4 without any Cl− ion, representing two kinds of passive potential regions which consisted of the primary low current density plateau and the secondary high current density plateau [13]. As Cr content increased, the primary passive current density decreased, exhibiting the secondary passive region disappeared. However, in the case of pure Fe, the high current density region was
Fig. 1. Apparatus for abrading electrode used in the study of repassivation kinetics of free electrode surface. during testing. A platinum wire and a saturated calomel electrode (SCE) were used as the counter and reference electrodes, respectively. Potentiodynamic polarization experiments were performed on specimens using a CMS 105B potentiostat controlled by a computer software (CMS 105). An abrading electrode technique was employed to obtain information on the repassivation kinetics of oxide films repassivated on the surface of specimens after exposure of the bare surface to the solution. The electrode was abraded with an SiC disc while immersed in the electrolyte to acquire current transients as shown in Fig. 1. The specimen was mounted in the center of a polyethylene resin (diameter = 1 cm) with a exposed area of 8 × 10−2 cm2 which was welded to a current collector of copper wire. The specimen surface was renewed by abrading with the SiC disc sintered from 1200 grit. The disc was fixed to the rotating shaft operated by dc motor. A saturated calomel electrode (SCE) was used as the reference electrode, located nearest to the working electrode. Finally a graphite was used as the counter electrode. During the abrading, the specimen was instantaneously raised by a mechanical spring. The resulting current transient was acquired at a constant applied potential from the moment just after interrupting the abrading action on the specimen. In order to evaluate the electrochemical behavior of repassivated films, ac impedance spectroscopy measurements were performed by using a potentiostat and a lock-in amplifier (SR810) with an electrochemical impedance software (CMS300) under static conditions upon completion of the abrading tests. Measurements were taken by superimposing an ac voltage 10 mV amplitude on the measured rest potential of the abraded specimen with the frequency range from 0.01 Hz to 10 kHz.
3. Results and discussion 3.1. Effect of Cr content on the passivity and the repassivation Stainless steels basically derive their passive characteristics from alloying with Cr. Alloying element, Cr is known to increase resistance to pitting by the increase of the pitting potential in the noble direction, reducing the passive current density. This is illustrated clearly by the data for a series of Fe–Cr–Ni alloys in a sulfuric acid solution [10], and for Fe–Cr alloys in a chloride solution [11], respectively. A.J. Sedriks summarized the effect of alloying elements in stainless steels on the anodic polarization curve [12]. According to his schematic summary, Cr increases
Fig. 2. Potentiodynamic polarization curves for Fe–Cr alloys in deaerated 0.1N H2 SO4 without and with Cl− ion according to the variation of Cr content: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.1N NaCl.
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Fig. 3. Decay of current density with time during repassivation of Fe–Cr alloys at the applied potential of −0.25 V in deaerated 0.1N H2 SO4 .
observed up to +900 mV above which a steep decrease in current followed by a low current density region was noted. Since this region corresponds to the secondary passive potential region of the anodic curves of Fe–Cr alloys, in addition to Cr oxide, the formation of Fe oxide seems to play some role on the secondary passive film of Fe–Cr alloys. However, since the primary passive potential region of Fe–Cr alloys agrees with the passive potential region of pure Cr and furthermore, with increasing Cr content, the polarization behaviors of Fe–Cr alloys approximates that of pure Cr, the Cr oxide on the surface of specimen plays the most important role on the formation of primary passive film. In order to investigate the effect of Cl− ion on the passive film, potentiodynamic curves for Fe–Cr alloys in aqueous 0.1N H2 SO4 with 0.1N NaCl is shown in Fig. 2(b). Pure Cr showed the similar kinds of polarization behaviors irrespective of Cl− ion concentrations while pure Fe did not display the secondary passive region due to addition of Cl− ion, representing Cr oxide film has the more resistance to corrosion than Fe oxide film in the chloride ion containing solutions. Therefore, based on the results of polarization curves, the passivity of Fe–Cr alloys may be mainly attributed to Cr oxide film whereas Fe oxide film has the less influence on the passivation behavior in the chloride ion containing acidic solutions The repassivation kinetics of Fe–Cr alloys can be expressed as follows. i = At −n
(1)
where i denotes the anodic current density, A is the constant, t is the time, and n is the repassivation rate parameter representing the slope of log i − log t plot. Based on Eq. (1), the value of n can be considered to be an indirect measure of the growth of the oxide on the bare surface as a function of applied anodic potential. Fig. 3 shows the potentiostatic current density decays plotted in log i versus log t obtained from the moment just after interrupting the abrading for Fe–Cr alloys applied at −0.25 V
in deaerated 0.1N H2 SO4 . Repassivation current density, i with time, t was observed to obey the linear relationship of log i versus log t [14–16]. The anodic current density of pure Fe was nearly constant maintaining high values with time on a logarithmic scale while that of pure Cr showed the lowest among pure Fe, pure Cr and Fe–Cr alloys. As the Cr content increased in Fe–Cr alloys, repassivation current densities decreased, slope −n of log i = k − n log t approaching to that of pure Cr. Pure Fe did not show any repassivation behavior at −0.25 V, in deaerated 0.1N H2 SO4 . These demonstrate that alloying element Cr in Fe–Cr alloys plays an important role in the rapid repassivation after fracturing the passive film Fig. 4 shows the decay of the current density with time in logarithmic scale during the repassivation of pure Cr with varying applied potential and Cl− ion concentration in deaerated 0.1N H2 SO4 solution. In the case of 0.1N H2 SO4 solution without Cl− ion, regardless of applied potential, pure Cr samples were completely repassivated with increasing time, showing that repassivation current densities increased with increasing applied potential in the passive potential range (Fig. 4(a)). However, in 0.1N H2 SO4 solution containing 0.1N Cl− ion, repassivation did not maintain at applied potential of 0.6 V, showing n = 0 after t = 200 s (Fig. 4(b)). As shown in Fig. 4(c), when the chloride ion concentration was increased to 0.6N NaCl and the anodic potential of 0.3 V was applied, an ascending anodic current density was observed, after t = 20 s, indicating the growth of stable pits [17]. Thus, in the case of absence of chloride ion, pure Cr showed the same repassivation kinetics irrespective of applied potential but in the presence of the chloride ion, even pure Cr was not able to become completely repassivated, as chloride ion concentration and applied potentials increased. In order to investigate effects of Cr content and applied potential on the corrosion resistance, ac impedance spectra for pure Cr and Fe–Cr alloys were measured at the open circuit potential in deaerated 0.1 M Na2 SO4 solution just after repassivation during 200 s. at various applied potentials −0.25, 0.3 and 0.6 V in deaerated 0.1N H2 SO4 solution. The passive oxide films formed by repassivation were analyzed by using the equivalent circuit model [18] (Fig. 5), which admits a one time constant Warburg diffusion model and the non-linear least squares (NLLS) fitting method. The charge transfer resistance (Rct ) of the repassivated film tended to increase with increasing Cr content and applied potential (Fig. 6(a)) while the capacitance of passive film (Cox ) tended to decrease with the increase in Cr content (Fig. 6(b)). The present results on the improvement of repassivated film with increasing Cr content may be explained by the investigation on the role of Cr on the passive film of Fe–Cr alloys in sulfuric acid carried out by J.A.L. Dobbelaar et al. [19]. They found that the porosity in a passive film decreases with increasing Cr content and/or applied potential, improving the protective quality of the film. 3.2. Effect of Mo content on the passivity and the repassivation Fig. 7 shows potentiodynamic polarization curves for Fe–18Cr–xMo (x = 1, 4, and 6 wt.%) alloys in aqueous 0.1N
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Fig. 4. Decay of current density with time for during repassivation of pure Cr at varying applied potential and Cl− concentration in deaerated 0.1N H2 SO4 solution: (a) 0.1N H2 SO4 , (b) 0.1N H2 SO4 + 0.1N NaCl and (c) 0.1N H2 SO4 + 0.6N NaCl.
H2 SO4 without and with Cl− ion concentrations of 0.1 N. In aqueous 0.1N H2 SO4 without Cl− ion, the increase in Mo content showed approximately same anodic polarization curves, compared with that of Fe–18Cr alloy without any Mo (Fig. 7(a)). However, in aqueous 0.1N H2 SO4 with Cl− ion, as Mo content increased, the pitting potential was raised, representing the passive potential region was extended in width (Fig. 7(b)). At the same time, the critical current density for passivation, ic decreased with increasing Mo content [20]. This suggests that the higher Mo content promotes the easier passivation and the more stable passivating oxide film. According to the results of Auger Electron Spectroscopy (AES) by Ogawa et al. [21] and X-ray Photo-electron Spectroscopy (XPS) by Hashimoto et al. [22], Mo can be interpreted
to play a significant role not only to promote the increment of the Cr content in the passive film but also to lead to the formation of stable and homogeneous passive films due to the decrease in the activity of surface sites. Montemor et al. [23] made analytical experiments for the passive film shown as a duplex structure for the oxides in stainless steels. Based on their results for AES and XPS, Mo was observed to be present in the internal Cr rich layers as well as at the Cr oxide/metal interface, resulting in an enrichment of Cr with the assistance of Mo in the internal region of the passive film. They also suggested that Mo introduces changes in the properties and the defect structure of the oxide film. Those experimental results associated with both the increase in pitting potentials and the extension of passive potential region in width were observed in Fig. 7, as Mo content increased.
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slope is in good agreement with minute pits on the metal surface observed after abrading Fe–18Cr alloy at −0.25 V in deaerated 0.1N H2 SO4 + 0.6N NaCl solution. In addition, Fig. 9(a) and (b) shows that the increase in Mo content causes the faster repassivation by reducing peak current density transients regardless of the variation of repassivation rate. Fig. 10 shows impedance measurement results in Nyquist plots for Fe–18Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys carried out just after performing abrading electrode test at an applied potential of −0.25 V. The ac impedance presented here were measured at the open circuit potential, Eoc in 0.1N H2 SO4 and 0.1N H2 SO4 + 0.6N NaCl, respectively. The diameter of the semicircle i.e. the charge transfer resistance, Rct increases with increasing Mo content. The increase in Rct with Mo content may be interpreted in the fact that the presence of Mo in repassivated film may shift the alloy from the active region to the pre-passive region, and form a pre-cursor film for passivation, since Rct of Fe–18Cr alloy without Mo is quite low. J.H.
Fig. 5. Equivalent circuit model for simulation of Fe–Cr alloys [18].
Fig. 8 shows the current density transients of Fe–18Cr and Fe–18Cr–xMo (x = 1, 4, 6 wt.%) alloys with time carried out just after performing abrading electrode test in deaerated 0.1N H2 SO4 solution at an applied potential of −0.25 V. Current density transient peak values decreased as Mo content increased, representing that addition of Mo leads to a significant decrease in the transient heights. Fig. 9 shows the potentiostatic current density decays plotted in log i versus log t obtained from the moment just after interrupting the abrasion of Fe–18Cr–xMo alloys applied at −0.25 V in deaerated 0.1N H2 SO4 (Fig. 9(a)) and 0.1N H2 SO4 + 0.6N NaCl (Fig. 9(b)). Repassivation current density with time was observed to obey the linear relationship of log i versus log t. In deaerated 0.1N H2 SO4 both Fe–18Cr and Fe–18Cr–xMo (x = 1,4, and 6 wt.%) alloys showed all anodic current densities continuously decreased with time in logarithmic scale. However in deaerated 0.1N H2 SO4 + 0.6N NaCl, anodic current density of Fe–18Cr alloy descended linearly with time in logarithmic scale and then the value of current density approached to nearly constant while those of Fe-18Cr–xMo alloys continuously descended linearly with time in logarithmic scale. The fall of anodic current density can be explained by the fact that the repassivation rate of passivating oxide film dominated over its dissolution rate [24]. The slope, −n of log i versus log t close to −1 is typical for passive films growing without the superimposed dissolution of passive films while 0 or positive slopes represent uniform corrosion or pitting corrosion, respectively. Therefore, the variation of a slope of Fe–18Cr with time, at −0.25 V in deaerated 0.1N H2 SO4 + 0.6N NaCl, imply the possibility of the dissolution of repassivated layers of Fe–18Cr after around 100 s whereas slopes of Fe–18Cr–xMo alloys close to −1 with time indicate the formation of the sound and protective film without the dissolution of passive layers during abrading tests. This possibility of pitting corrosion due to the variation of
Fig. 6. Changes of Rct and Cox of Fe–Cr alloys obtained from the impedance spectra with varying applied potentials and Cr content: (a) Rct and (b) Cox .
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Fig. 8. Current density transients curve of Fe–18Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys in deaerated 0.1N H2 SO4 solution at an applied potential of −0.25 V.
and stabilization of Cr oxide film [16]. It has also been suggested that Mo retards the corrosion process by adsorption [30], by formation of Mo compounds [31], by synergistic interaction of Mo ions with other oxides of the passive film [20,32] and by elimination of the active surface sites through formation of Mo oxides [22]. Even though there are analytical results showing the presence of Mo, in other studies, Mo was not found in the passive film [22,27,28]. Therefore, the satisfactory explanation for the role of Mo on corrosion resistance is still a subject of discussion. Based on these present experimental results, the increase in Mo content raises the pitting potentials, producing a more stable passive film and increasing the charge transfer resistance. Furthermore, the addition of Mo facilitates the faster repassivation by reducing peak current density transients,
Fig. 7. Potentiodynamic polarization curves for Fe–18Cr–xMo alloys in deaerated 0.1N H2 SO4 without and with Cl− ion according to the variation of Mo content: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.1N NaCl.
Gerretsen and J.H.W. de Wit investigated the role of Mo on the passivation of Fe–Cr alloys [25]. According to their results, Mo increases the dissolution rate of the alloys, but the increase is compensated by the faster formation of the passive film, thus increase the resistance to corrosion of the Fe–Cr alloys, which is in agreement with the present experimental results showing that the addition of Mo to Fe–18Cr causes the faster repassivation during abrading and the higher charge transfer resistance measured by ac impedance spectroscopy after abrading test. The beneficial effect of Mo in the corrosion resistance of stainless steels has been known to be attributed to the following several factors. These include an enrichment of Mo in the passive film [20,26] or just below the passive film [27,28], enrichment of Cr in the oxide layer [28,29], thickening of the passive film
3.3. Effect of N content on the passivity and the repassivation Fig. 11 shows potentiodynamic polarization curves for type 316L (0 wt.% N), type 316LN (0.042 wt.% N), type 316LN (0.1 wt.% N) and type 316LN (0.15 wt.% N) austenitic stainless steels in deaerated 0.1N H2 SO4 with varying Cl− ion concentration. As shown in Fig. 11(a), N addition to type 316LN in deaerated 0.1N H2 SO4 without Cl− ion did not make any noticeable difference except for somewhat reducing the critical current density for passivation, ic . Even though Cl− ion concentration increases up to 0.6N, the addition of N seems to have the anodic polarization curves of type 316LN stainless steels unchanged. However, the anodic polarization curve of type 316L without N in deaerated 0.1N H2 SO4 + 0.6N NaCl solution, shows the much higher passive current density, ipass and critical anodic current density for passivation, ic than type 316LN with N, indicating that N addition promotes the easier passivation and the more stable passivating oxide film. In the similar manner of Mo, the
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linearly with time in logarithmic scale but once the current densities reached minimum current values, the current densities were observed to ascend. This ascending anodic current density can be attributed to the formation and growth of pits due to the passive film breakdown, depending upon applied anodic potential and chloride ion concentration [17]. However the anodic current densities of type 316LN (0.042 wt.% N) and type 316LN (0.15 wt.% N) continuously descended linearly with time in logarithmic scale irrespective of applied potentials, indicating that the repassivation rate of passivating oxide film dominated over its dissolution rate on the bare surface. The role of N in corrosion resistance of stainless steels can be explained by suppressing the dissolution of oxide film and by facilitating the repassivation of the passive film, which cause a decrease in the height of current transients.
Fig. 9. Decay of anodic current density with time for repassivation of Fe–Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys at −0.25 V in: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.6N NaCl.
addition of N leads to increase the pitting potentials, widen the passive potential region, and reduce the critical anodic current density, promoting the faster formation of passive films and a more stable passive film. These results seems to be related with those reported by Kamachi Mudali et al. [33] that pit initiation was delayed and pit growth was significantly decreased with the addition of N to pure iron. Fig. 12 shows the potentiostatic current density decays plotted in log i versus log t for type 316L (0 wt.% N), type 316LN (0.042 wt.% N) and type 316LN (0.15 wt.% N) austenitic stainless steels obtained at applied potentials of 0, 0.4, and 0.6 V in deaerated 0.1N H2 SO4 + 0.1N NaCl. Repassivation current densities of type 316L without N at 0.4 and 0.6 V initially descended
Fig. 10. Impedance spectra in Nyquist plot for Fe–18Cr and Fe–18Cr–xMo (x = 1, 4 and 6 wt.%) alloys carried out at Eoc just after performing abrading electrode test at an applied potential of −0.25 V in: (a) 0.1N H2 SO4 and (b) 0.1N H2 SO4 + 0.6N NaCl.
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Fig. 11. Potentiodynamic polarization curves for 316L and 316LN with different N content in deaerated 0.1N H2 SO4 with varying Cl− ion concentration: (a) 0.1N H2 SO4 , (b) 0.1N H2 SO4 + 0.1 NaCl and (c) 0.1N H2 SO4 + 0.6N NaCl.
Fig. 13 shows impedance spectra in Bode plot for type 316L (0 wt.% N), type 316LN (0.042 wt.% N), and type 316LN (0.15 wt.% N) carried out at the open circuit potential, Eoc after performing potentiostatic current density transient measurements at applied potentials of 0 and 0.6 V in deaerated 0.1N H2 SO4 . The charge transfer resistance, Rct of type 316LN increased with N addition by 10 times compared with that of type 316L without N. Fig. 13(b) shows charge transfer resistance of repassivated film increased with the increase in applied potential from 0 to 0.6 V in deaerated 0.1N H2 SO4 . It is suggested that the passive film formed at higher applied potential is more protective than that produced at lower applied potential within the
passive potential region of deaerated 0.1N H2 SO4 solution. This indicates that the higher potentials apply in the passive potential region, the more stable oxide films form in deaerated sulfuric acid solution. The influence of N on the corrosion resistance of stainless steel has received much attention in recent years. There is a general consensus that N additions improve the resistance to pitting corrosion, increasing the pitting potential in aqueous chloride solutions. The present experimental work shows that N also accelerate repassivation. In terms of the role of N in the resistance to corrosion, Osozawa and Okada [34] attributed the improved pitting
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Fig. 12. Decay of anodic current densities with time for repassivation kinetics of 316L, 316LN (0.042 wt.% N) and 316LN (0.15 wt.% N) with various applied potentials in deaerated 0.1N H2 SO4 + 0.1N NaCl solution: (a) 316L; (b) 316LN (0.042 wt.% N) and (c) 316LN (0.15 wt.% N).
resistance of N-containing stainless steels to the formation of ammonium ions at pit sites, which increased the pH. Newman and Shahrabi [35] classified the effects of N on pitting corrosion resistance into three steps, namely (1) consumption of acid in pit nuclei by nitrogen dissolution, i.e. N + 4H+ + 3e → NH4 + , (2) enrichment of N on the passivated surface, and (3) enrichment of N on active surfaces. Newman et al. [36] also observed enrichment of Mo and N at the passive film/metal interface and attributed this enrichment to the prevention of the dissolution of the substrate following destruction of the passive film. JargeliusPettersson [37] found the N additions increased the critical temperature for pitting corrosion, indicating a synergistic interaction between Mo and Olefjord and Wegrelius [38] observed
that the high N-containing steel was repassivated while the pits were growing on the low N-containing steel, suggesting a model for a synergistic effect between Mo and N. However, only a few studies concerning the effect of Mo and N on the repassivation kinetics has been reported, and it is proposed that N-containing stainless steels without Mo, did not show any influence of the N content on the polarization behavior [39,40]. Therefore, even though based upon the present studies, alloying elements, Mo and N promote the passivation by formation of more stable passive film and the repassivation by increasing the formation rate of passive film, the exact mechanisms of the role of Mo and N on the repassivation kinetics and the exact interaction between N and Mo remains still to be clarified.
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with increasing applied potential within the passive potential region. 2. The increase in Mo content raises the pitting potentials and reduces the critical anodic current density for passivation in aqueous 0.1N H2 SO4 with Cl− ion. The addition of Mo facilitates passivation, produces the more stable passive film, causes the faster repassivation by reducing peak current density transients and finally promotes the resistance to corrosion by increasing the charge transfer resistance. 3. Compared type 316L and type 316LN stainless steels, type 316LN to which N is added shows a much lower passive current density and critical anodic current density for passivation. In the similar manner of the addition of Mo, the addition of N increases the repassivation rate of the oxide film, reducing the corrosion rate. Acknowledgment The author gratefully acknowledges the financial support by Korea Science and Engineering Foundation (KOSEF). References [1] [2] [3] [4] [5]
[6]
Fig. 13. Impedance spectra in Bode plot for 316L, 316LN (0.042 wt.% N) and 316LN (0.15 wt.% N) carried out at Eoc , after performing potentiostatic current density transient measurements in deaerated 0.1N H2 SO4 at applied potentials of (a) 0 V and (b) 0.6 V.
[7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
4. Conclusions 1. The corrosion potentials and the passive current densities of Fe–Cr alloys were decreased with the increase in Cr content. The passivity of Fe–Cr stainless steels was mainly attributed to primary Cr oxide film. In addition to Cr oxide film, Fe oxide may have something to do with secondary passive film. Addition of Cl− ion to 0.1N H2 SO4 solution resulted in the contracted passive potential region and the increased passive current density. The increase in Cr content of Fe–Cr alloys increased both the repassivation rate and charge transfer resistance, Rct . In the case of addition of Cl− ion, the tendency of suppressing repassivation rate increased
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