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Electrochemical characterization of oxide formed on chromium containing mild steel alloys in LiOH medium
Materials Chemistry and Physics 145 (2014) 499e509
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Electrochemical characterization of oxide formed on chromium containing mild steel alloys in LiOH medium Veena Subramanian, Sinu Chandran, H. Subramanian, P. Chandramohan, S. Bera, S. Rangarajan*, S.V. Narasimhan Water and Steam Chemistry Division, Bhabha Atomic Research Centre Facilities, Kalpakkam 603102, Tamilnadu, India
h i g h l i g h t s
g r a p h i c a l a b s t r a c t
High temperature oxides formed on Cr containing mild steels are less defective. Defect densities of oxides decrease with increase in Cr content in the alloy. O2 in solution greatly influences the nature and defect chemistry of oxides.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 June 2013 Received in revised form 10 January 2014 Accepted 1 March 2014
Flow accelerated corrosion leads to wall thinning of outlet-feeder pipes in the primary heat transport system of pressurized heavy water reactors and can even necessitate enmasse feeder replacement. Replacement of carbon steel 106-grade-B (CS) with chromium containing carbon steel reduces the risk of this failure. This paper discusses the role of small additions of chromium in modifying the properties of the oxide film. CS and chromium containing mild steels viz., A333, 2.25Cre1Mo and modified 9Cre1Mo alloy were exposed to primary heat transport (PHT) system chemistry conditions. The oxide films formed were characterized by electrochemical and surface characterization techniques. MotteSchottky analysis showed donor type of defects. The densities of defects in the oxides of chromium containing alloys were 3e15 times less than that in CS. In presence of w200 ppb of dissolved oxygen, the oxides formed were hematite with two orders of magnitude smaller concentration of defects as compared to that formed under reducing conditions. These results suggest that the presence of chromium lowers the defect density of the oxide film and thus ensures a reduced corrosion rate. Ó 2014 Elsevier B.V. All rights reserved.
Keywords: A. Alloys C. Electrochemical techniques C. X-ray photo-emission spectroscopy (XPS) D. Corrosion
1. Introduction In pressurized Heavy Water Reactors (PHWR), the corrosion of the reactor primary circuit materials is minimized by maintaining an alkaline pH (at 25 C) of 10.2e10.6, by the addition of LiOH. Despite maintaining this stringent water chemistry condition, some of the CANDU type of reactors and PHWRs are facing the * Corresponding author. Tel.: þ91 44 27480203; fax: þ91 44 27480097. E-mail addresses: [email protected] (V. Subramanian), [email protected]. in, [email protected] (S. Rangarajan). http://dx.doi.org/10.1016/j.matchemphys.2014.03.003 0254-0584/Ó 2014 Elsevier B.V. All rights reserved.
problem of outlet feeder thinning [1e3]. This phenomenon was attributed to flow accelerated corrosion (FAC). FAC is a process in which carbon steel piping and components corrode in the presence of flowing water or steamewater mixtures with low dissolved oxygen. As water flows over the carbon steel material, stable surface oxide layer (typically Fe3O4) dissolves into the flowing stream, thinning the walls of piping over time and resulting in failures due to rupturing [4e6]. The rate of FAC depends upon hydrodynamic factors such as flow velocity, pipe roughness and geometry of the flow path etc., environmental factors like temperature, pH, iron unsaturation, oxygen concentration and metallurgical factors like
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the composition of steel and residual stress. In addition to the general enhancement in corrosion rate due to high velocity prevailing in feeders, limited area of piping face further enhanced wear rate. This is attributable to pipe geometry which causes localized turbulence and eddies [7]. Hence this type of failure is specific only to outlet feeders of certain PHWRs and CANDUs and to certain channels. Periodic in service inspection campaigns are carried out in reactors to assess the residual life of these piping. However, exhaustive measurements in an operating plant are not always feasible and may even be hampered by access restrictions. So predictive models, which account for fluid dynamics and chemistry at each of the locations, are used to make reliable predictions and aid corrosion assessment of inaccessible areas. Many commercial software like CHECKWORKS, BRT CICERO, RAMEK, COMSY etc. are available to predict FAC [8]. The predictions are based on the assumption that the dominating factor is the presence of magnetite on the corroding surface, its solubility under the prevailing chemistry conditions, mass transfer of the dissolved ions to the fluid and subsequent mass transport to the bulk fluid. These models take in to account the fluid dynamics and chemistry at each of the susceptible locations and make reliable extrapolations for inaccessible areas. The input variables for these models are temperature, pipe geometry, fluid velocity, void fraction, chemical composition of the structural material, pH and dissolved oxygen. Research has shown that trace amounts of alloying elements e particularly chromium, greatly reduce the rate of FAC [9,10]. A study by Cheng et al., demonstrated that an increase in Cr content in the steel can reduce the corrosion rate of steel [11]. The passivity of chromium containing steels is attributed to the formation of a Cr2O3 film. The development of passivity due to incorporation of Cr in MeOeM network of an iron rich alloy requires that the chromium content should be atleast 10% [12e14]. The amount of chromium that is required to be present in the alloy to reduce the FAC rate has been experimentally found to be w0.04% [9,10]. Hence, it is of interest to know how such small amounts of chromium would influence the oxide dissolution and subsequent corrosion behavior. In this paper, an attempt is made to understand the role of chromium in modifying the properties of oxides formed on chromium containing mild steel alloys compared to that of Carbon Steel 106 grade B alloy (CS), which is the current feeder material in some of the Indian PHWRs. ASTM A 333 grade 6 alloy contains about 0.33% of chromium and is proposed as a new material to replace the CS feeders. Hence, this study compares the oxide properties of mild steel alloys viz., A333 (Cr w 0.33%), 2.25Cr1Mo alloy (Cr w 2.25%) and modified 9Cr1Mo (Cr w 9%) along with carbon steel 106 grade B (CS). Oxygen addition to the feedwater is practiced in some pressurized water reactors (PWRs) to protect the carbon steel piping from undergoing FAC, on the secondary side. The added oxygen reacts with the magnetite and converts it to hematite, which has four orders of magnitude lower solubility compared to magnetite. The effectiveness of this modified oxide as a barrier to FAC should also have its origin in the defect density and electronic properties of the oxide film. It is desirable to compare the extent of protection offered by chromium in the alloy and the beneficial effect of added oxygen in the process fluid. This paper also discusses the effect of dissolved oxygen for different chromium containing alloys. The passive films formed at high temperature are semiconducting in nature and the metal release from these oxide films depends on the number density and diffusivity of the ionic point defects [15e18]. Typical donor defects in oxides are oxygen vacancies and cation interstitials and the acceptor defects are mainly cation vacancies. The influence of chromium in changing the defect structure of the oxide was studied by MotteSchottky analysis of the measured capacitance as a function of applied electrode potential.
Table 1 Compositions of the alloys studied.
C Cr Si Ni Mn Mo Cu Co Sn Fe a
Carbon steel
A-333
2.25Cr1Mo
Modified 9Cr1Mo
0.1 <0.3 0.28 <0.14 0.75 <0.03 0.16 <0.03 0.015 Bal
0.24 <0.3 (0.5)a 0.42 <0.14 1.07 e 0.1 <0.013 0.011 Bal
0.065 2.2 <0.2 0.15 0.5 1 0.055 <0.009 <0.008 Bal
0.095 8.1 0.4 0.13 0.38 0.9 0.07 0.025 <0.008 Bal
Measured by atomic absorption spectrometry.
The stability of the film in LiOH and the changes occurring at the oxide/solution interface were studied by electrochemical methods such as Electrochemical Impedance Spectroscopy (EIS) and Potentiodynamic Anodic Polarization (PDAP). 2. Experimental All the chemicals used were GR/AR grade. LiOH was prepared by passing Li2CO3 through OH form of resin. The outlet solution contained 1.5 ppm of Li as analyzed by flame photometry and the pH and conductivity were found to be 10.2 and 45 mS cm1 at 25 C respectively. The three different chromium containing alloys and carbon steel (coupons of approximately 1.2 cm 1.2 cm 0.4 cm) were exposed to lithiated water (LiOH) at 245 C in a static autoclave so as to form their respective oxides for 20 days. The surfaces of the specimens were ground up to 800 grit by using silicon carbide papers, cleaned with ultra-pure water and finally with acetone before exposure. The solution was purged with argon gas for 3e4 h before raising the temperature to ensure removal of dissolved oxygen. One of the parameter that can influence the nature of oxide film is the oxygen content in the solution. Our previous study, wherein the oxide films were formed on these alloys at high temperature by allowing oxygen but only for a short duration, resulted in highly resistive oxides. The composition of the oxides on all the alloys were found to be a mixture of magnetite and hematite in the ratio 3:1 [19]. In the present study, a recirculating autoclave facility was used in order to maintain a steady concentration of oxygen throughout the experiment. In this set of experiments, the specimens were exposed to high temperature simulated PHT conditions, of 245 C and pH25 of 10.2 (by LiOH) in a 2.4 L re-circulating autoclave for 16 days. The flow rate was 10 LPH, thus the water was renewed 4.2 times in an hour. Some of the advantages of switching over to recirculating system from static system were that the solution could always be regenerated and DO could be maintained at 200 10. The chemistry parameters in the source tank, like pH, specific conductivity, dissolved oxygen and redox potential could be measured using an online low temperature chemistry monitoring system. In addition, corrosion potential could be measured online at high temperature and iron-
Table 2 Estimates of oxide thickness on the alloys. Material
Oxide thickness in mm after 20 d exposure to 245 C (deaerated condition)
Oxide thickness in mm after 16 d exposure to 245 C (with DO)
CS A-333 2.25Cr-1Mo Mod. 9Cr1Mo
0.77 0.81 1.0 1.09
1.1 0.91 0.66 0.21
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Fig. 1. SEM micrographs of oxide formed on alloys exposed to deaerated conditions a) CS b) A-333, c) 2.25Cr-1Mo and d) 9Cr-1Mo.
unsaturation could be achieved with the purification flow through a lithiated mixed bed resin at the downstream of the autoclave. Iron-unsaturation in coolant is one of the chemistry parameters which is known to aggravate FAC [20]. The electrochemical potential measured for CS and A333 increased gradually from 0.4 V to 0.1 V and remained steady at 0.1 V after 70 h of exposure. The increase in potential was due to the formation of a stable oxide film. Electrochemical and surface characterization studies of these oxide coated specimens were carried out at room temperature. For
electrochemical studies, the oxide coated specimen (area exposed w1 cm2) was made as the working electrode. The electrolyte used was LiOH, prepared as described above. The electrochemical measurements were made using AUTOLAB PGSTAT 30. Three-electrode configuration was used with a platinum foil serving as the counter electrode and a saturated calomel electrode (SCE) as the reference electrode which was connected through a luggin capillary. In this paper, for the room temperature studies, the values of the potentials are given with respect to SCE only. All the experiments were carried out under deaerated condition by purging the solution with
Fig. 2. SEM micrographs of oxide formed on alloys exposed to oxygenated conditions a) CS b) A 333, c) 2.25Cr-1Mo and d) 9Cr-1Mo.
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Fig. 3. Raman spectra of oxide formed on specimens exposed under deaerated conditions.
argon. A scan rate of 0.5 mV s1 was used for potentiodynamic polarization measurements. Electrochemical impedance spectra were obtained by applying an AC voltage (10 mV amplitude) to the sample at open circuit potential (OCP) in the frequency range of 10 kHze10 mHz. Capacitance measurements on these oxidecoated specimens were carried out from 1.2 V to 0.8 V range and MotteSchottky analyses were performed to assess their defect densities. Scanning electron micrographs (SEM) were taken by using a PHILIPS ESEM XL-30 system. Raman spectra were recorded by using micro Raman Spectrometer (LABRAM HR800) equipped with 514 nm argon ion laser. Laser was focused using 100 objective lens on sample surface with a laser power of 0.5 mW. A grating with 1800 groves/mm was used with a 30 s data acquisition time. X-ray Photoelectron Spectroscopic (XPS) studies were done using VG ESCA Lab MK 200 equipment. The compositions of the alloys were estimated by using direct reading optical emission spectroscopy and are given in Table 1.
3. Results and discussion 3.1. Oxide thickness measurement Estimates of the thickness of the oxide films were made by descaling them using modified Clarke’s solution [21]. Because of its chromium content, oxides on 2.25Cre1Mo and 9Cre1Mo could be descaled only partially using this solution. Hence, a two step process involving caustic permanganate solution (10% NaOH, 4% KMnO4 at 95 C) followed by a treatment with ammonium citrate (10%) was used [22]. The estimates of the thickness of oxides formed on the four alloys, in the presence and absence of dissolved oxygen, are given in Table 2. 3.2. Surface characterization studies 3.2.1. Scanning electron microscopy Scanning electron micrographs of the oxide coated specimens, after exposure to LiOH at 245 C for 20 days under deaerated
Fig. 4. Raman spectra of oxide formed on specimens exposed to controlled oxygenated conditions.
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Fig. 5. XPS spectra of oxide formed at controlled oxygen conditions: (a) Fe 2p and (b) Cr 2p.
conditions, are presented in Fig. 1. The films were made up of crystalline oxide particles and their sizes were found to increase with time of exposure. With all the materials, particularly on the 20 days-exposed specimens, both smaller and bigger oxide crystals were observed, the latter being precipitated from the solution. Energy dispersive X-ray (EDX) microanalysis of these films confirmed that the major constituent was iron. The chromium contents of the films were found to be 0.3%, 0.2%e0.6%, 3%, 6e9% on CS, A-333, 2.25Cre1Mo and 9Cre1Mo alloys respectively. Increase in exposure time did not result in appreciable change in oxide composition for all the materials tested and they are not discussed further in this paper. SEM micrographs of the specimens exposed in the presence of controlled oxygen environment are given in Fig. 2. The surface morphologies of the crystallites were quite different; the typical octahedral edges were not seen as observed on the specimens exposed to deaerated conditions. EDX analysis indicated only the presence of iron and oxygen on the particles. Chromium was seen only on 2.25Cr1Mo and 9Cr1-Mo and the measured concentrations were almost the same as on the base metal. Since the depth of penetration is of the order of w0.5 mm in EDX, the base metal contribution to the chromium content observed in EDX spectra cannot be ruled out.
3.2.2. Raman spectroscopy Raman spectra for the specimens exposed to reducing conditions are given in Fig. 3. The principle peak at 670 cm1 (A1g) is typical of magnetite phase as observed for the oxide formed on CS and A333. On 2.25Cr-1Mo and 9Cr-1Mo, the incorporation of Cr into the ferrite lattice has resulted in a shift of the peak to higher wave numbers. Magnetite is an inverse spinel with all Fe2þ occupying the octahedral (Oh) sites, half of the Fe3þ occupying the tetrahedral (Td) and the remaining Fe3þ occupying Oh sites. It can be represented as [Fe3þ]Td[Fe2þ, Fe3þ]Oh. Here, the Raman active mode is the FeeO stretch involving the [Fe3þ]Td along with the neighboring four oxygen atoms. The strength of the A1g vibrational mode is very sensitive for any substitutions in magnetite. It is reported that the divalent ions like Zn2þ could occupy the Td sites in the magnetite lattice and affect the main peak at w670 cm1 to a greater extent and shift its position significantly [23,24]. However, in Cr3þ substitution, the ion will go and replace the Fe3þ in Oh site, and hence the change brought about in the main Raman active mode will be an indirect one. Hence, the shift could be marginal as observed in the present case. Under controlled oxygenated conditions, all the specimens developed a grayish black oxide film. In their Raman spectra, a
Fig. 6. XPS spectra of oxide formed at deaerated conditions: (a) Fe 2p and (b) Cr 2p.
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distinct and intense peak characteristic of a-Fe2O3 was observed at 227, 248, 294, 412, 500, 613 and 1322 cm1 (Fig. 4). The presence of the oxygen in the water, favored the formation of hematite as an outer layer in all the specimens. In the case of 9Cr1Mo, a small peak at w663 cm1 was seen corresponding to spinel or ferrite/chromite phase.
impedance due to bulk oxide and impedance due to charge transfer processes at oxide/solution interface. Oxide film here can be considered to be made of two distinct layers, the inner film whose composition depends only on the base metal and the outer film which forms due to the precipitation from the solution and hence can vary depending on the additives in the solution. Under the deaerated conditions, the inner and outer oxides are indistinguishable hence
3.2.3. X-ray photoelectron spectroscopy XPS studies on the oxides formed at controlled oxygen conditions confirmed the presence of chromium only on 9Cr-1Mo and iron was present in Fe3þ state on all the specimens (Fig. 5). The Fe 2p3/2 peak was found to be associated with a distinct satellite peak at around 719 eV (marked by an arrow) having a narrow peak width (3.5 eV) which indicated the formation of Fe2O3. For specimens exposed to deaerated conditions, the Fe 2p3/2 peak was broad and the satellite peak characteristic to Fe3þ (around 8 eV left to main peak) was not seen (Fig. 6(a)). In addition, the peak width was around 4.6 eV. All these observations indicated the presence of Fe2þ along with Fe3þ. Under both the conditions, Cr was observed only on 9Cr-1Mo. In both the cases, the Cr 2p3/2 peak was seen at 576.8 eV which indicated the presence of Cr3þ [25]. These observations from Raman and XPS suggested the formation of spinels with Cr. The atomic concentrations of Fe and Cr were evaluated from the peak areas and were found to be 70 at% and 30 at% respectively. Preferential dissolution of Fe, under deareated conditions, probably would have increased the Cr content in the oxide film. Similar observations were made with carbon steel, A333 and 2.25Cr-1Mo specimens. Thus, under deaerated conditions, the surface was dominantly composed of spinels. 3.3. Electrochemical characterization The impedance data was validated using the ‘linear Kramerse Kronig transformation’ function of the commercial software available with the frequency response analysis (FRA) unit from Autolab. Linear KramerseKronig test is based on an approach developed by M.E. Orazem and co-workers and further modified by Boukamp [26]. The matching of the KK transformed and the experimentally obtained data confirmed that the data satisfied all the criteria of a valid impedance data. Electrochemical impedance spectra were measured for all the oxide coated specimens that were exposed to deaerated and oxygenated conditions. The alloys were found to develop a duplex oxide film on their surfaces (Fig. 7). The inner film was found to have smaller particles (w0.1 mm) characteristic of their formation due to the oxidation of the base metal. The outer oxide had relatively bigger particles (1 mm) suggesting their formation due to precipitation from the solution. The outer oxide film could grow with the time of exposure. Hence, the physical model for analyzing EIS data can be represented by a metal covered with a duplex oxide. 3.3.1. EIS for specimens exposed in static autoclave under deaerated conditions EIS data are given in Nyquist and Bode format in Fig. 8(a), (b) and (c) respectively. Analysis of the impedance data was done by fitting them into an electrical equivalent circuit. The total impedance of the system is due to three components. The solution impedance, the
Fig. 7. Schematic of oxide growth as a function of exposure time.
Fig. 8. (a) Nyquist and (b) &(c) Bode plots for oxide covered specimens exposed to deaerated condition.
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Fig. 9. Electrical equivalent circuit used for fitting EIS data for specimens exposed to (a) deaerated and (b) oxygenated conditions.
the impedance spectra was found to fit to electrical equivalent circuits with both two and three time constants. Since no significant differences were observed in the c2 value among the fitted data, the simpler circuit with two time constant (Fig. 9(a)) was chosen for this particular case. In Fig. 8, the lines represent the theoretical fit of the experimental data by using this particular equivalent circuit. The impedance response at high frequency corresponds to bulk oxide and at low frequencies to the charge transfer processes occurring at the outer oxide/solution interface. Table 3 gives the fitted impedance parameters. RS is the solution resistance, Rox, the resistance of the oxide film, Cox, the oxide capacitance and Rct, the charge transfer resistance for the processes that may be occurring during the measurement at oxide/solution interface. The Constant Phase Element (CPE), observed is of the order of millifarads and thus corresponds to diffused double layer at the oxide/solution interface. CPE is the non-ideal capacitance and represents a true capacitance if the value of the power term, n, is close to 1. If the exponent value is close 0.5, it indicates diffusion (called Warburg element) as the rate determining process. The value of n away from 1 is also interpreted to represent a nonhomogeneous surface. The bulk oxide properties were observed to be due to the base metal composition as the carbon steel and A333 behaved similarly. Similar trend in bulk oxide properties was observed among the 2.25Cr-1Mo and 9Cr-1Mo alloys also. The oxide capacitance was lower and the resistance higher in 9Cr-1Mo suggesting a larger separation of charges in its oxide. The more resistive interfaces were indicated with increasing Rct values with increasing Cr-content. 3.3.2. EIS for specimens exposed in recirculating loop with controlled dissolved oxygen (DO) The presence of w200 ppb of DO throughout the experiment was ensured in this case. The films formed on all the alloys were found to be hematite. The impedance plots for all the exposed specimens are given in Fig. 10. The inset in the figure gives the impedance spectra at highest frequencies. In this case, the data could not be fitted to the simpler circuit (Fig. 9(a)). In addition to the larger errors it also resulted in the complete loss of information of intermediate frequencies. Hence the equivalent circuit represented in Fig. 9(b) was used. It is quite reasonable since here, the presence of hematite introduces an inhomogeneity in the bulk oxide. It can be seen in Fig. 10 that the experimental data fits quite well with the fitted curve simulating the three time constant equivalent circuit. Table 3 Fitted impedance parameters for specimens exposed to deaerated conditions. Material
Rs (U cm2) Cox (nF cm2) Rox RCT CPE (F cm2) n (U cm2) (kU cm2)
CS 144 A-333 94 2.25Cr-1Mo 257 9Cr-1Mo 301
25.3 24.8 27 14.3
588 604 355 452
59.5 95 234 297
0.3e3 0.2e3 0.3e3 0.6e4
0.8 0.9 0.9 0.8
Table 4 gives the fitted parameters for the impedance data. Rs, Rox, Cox, CPEo/s and Rct have the same meaning as described earlier in Section 3.3.1, for deaerated specimens. In addition, Cpore and Rpore, represents the capacitance and resistance inside the pores at the inner oxide and outer oxide interface. The charge transfer resistances for all the specimens were found to be similar and is fixed at about w500 MU during fitting of the data. The high value suggests a highly resistive oxide/solution interface. For all the specimens, the capacitances were found to be quite similar. The diffused double layer capacitances Cpore and CPEo/s were found to be w2 orders of magnitude lower than CPEo/s at under the deaerataed conditions suggesting a thicker diffused double layer at the interfaces in the oxygenated conditions. This can be attributed to charge distribution in a highly resistive oxide. The oxide film resistances were found to be comparable for all the specimens. The pore resistances of CS and A333 were found be relatively higher than that of 2.25Cr-1Mo and 9Cr-1Mo alloys. A possible reasoning can be given as follows. At lower chromium concentration in the oxide, Cr3þ is likely to be present in the octahedral sites due to the favorable site preference energy [27]. Here, because of the presence of oxygen, the chromium from the outermost layer gets oxidized to chromate and dissolves. Any loss of Cr3þ from the oxide by oxidation to Cr6þ will result in formation of cation vacancies in octahedral sites. Further the charge will be balanced by oxidation of adjacent Fe2þ to Fe3þ. This increased cation vacancy defect density can result in increased transport of the metal ions through the oxide and progressive conversion to a less protective oxide rich in Fe3þ can increase the solubility of the oxide layer, both manifesting as a decreased resistance. Hence the chromium depletion renders the outer oxides on 2.25Cr-1Mo and 9Cr-1Mo alloys more inferior in terms of resistance as compared to CS and A333. 3.3.3. PDAP studies Potentiodynamic anodic scans for specimens previously exposed under deaerated and controlled DO conditions are given in Fig. 11(a) and (b) respectively. The passive current density measured in these scans corresponds to the steady state current and in effect, the steady state concentration of defects in the film. Table 5 gives the range of passive current densities (Ip) deduced from these polarization scans. Ip was observed to decrease as CS > A333 > 2.25Cr-1Mo > 9Cr-1Mo. The Open Circuit Potential (OCP) for hematite coated specimens was observed to be always lower as compared to magnetite coated ones. This shift could be due to the lowering of exchange current density at the cathodic sites at a more resistive oxide. The Ip for hematite films were of the order of nano-amperes as compared to magnetite films which were of the range of micro-amperes. In presence of DO, A333 alloy showed the lowest passive current. This observation essentially confirmed the high resistivity of the films formed with continuous presence of DO in the system.
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Fig. 11. PDAP curves for specimens exposed under (a) deaerated and (b) controlled oxygenated conditions.
3.3.4. Capacitance measurements The magnetite and hematite films formed on the specimens are semiconducting in nature. The metal release from these passive films depends on the concentration of point defects present in them. MotteSchottky analyses were done on the measured capacitance values to calculate defect densities for all the exposed specimens. This method involves the measurement of apparent capacitance as a function of potential under depletion condition where the following relation is obeyed [28,29]:
1 2 kT E E ¼ FB 2 eεr ε0 N e Csc
(1)
where
Fig. 10. (a)Nyquist and (b) and (c) Bode plots for oxide covered specimens exposed to controlled oxygen condition.
Csc ¼ capacitance of the space charge region εr ¼ dielectric constant of the oxide (semiconductor) ε0 ¼ permittivity of free space N ¼ donor density (for n-type)/acceptor density (for p type)
Table 4 Fitted impedance parameters for specimens exposed to oxygenated conditions. Material
Rs (kU cm2)
Cox (nF cm2)
Rox (kU cm2)
Rpore (kU cm2)
Cpore (mF cm2)
RCT (MU cm2)
CPE (F cm2)
n
CS A-333 2.25Cr-1Mo 9Cr-1Mo
0.9 0.9 0.6 0.5
9.9 14.0 11.7 6.4
1.3 1.8 1.1 1.6
2.7 2.8 1.5 1.7
9.1 6.7 10.6 26.7
500 500 500 500
0.2e4 0.3e4 0.3e4 0.5e4
0.8 0.8 0.8 0.8
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Table 5 Ip for the specimens exposed to deaerated and controlled DO conditions. Material
Deaerated
CS A-333 2.25Cr-1Mo 9Cr1-Mo
Controlled DO
OCP (V)
Ip (A cm2) 106
OCP (V)
Ip (A cm2) 109
0.258 0.17 0.212 0.128
1.04 0.47 0.34 0.12
0.402 0.308 0.286 0.14
8000e80 20e0.1 90e3 50e1
E ¼ applied potential EFB ¼ flat band potential e ¼ electronic charge The oxides formed on these alloys were found to follow Motte Schottky relationship at a certain potential range. Fig. 12 shows the typical MotteSchottky plots at 10 Hz for the alloys exposed to deaerated conditions. It was reasonable to take this frequency for measurement, as in our study, at >1 Hz, only the imaginary impedance of the system was found to respond to change in frequency. It is reported in literature that in polycrystalline semiconductor, the MotteSchottky behavior was attributed to the low ionic mobility in the depletion layer, which responds only at low frequencies [30,31]. It can be seen that, except for 9Cr-1Mo, the flat band potential is z0.5 V for all the alloys. Oxides on all the materials show n-type behavior at potentials noble to flat band potential. The donor type defects here are O2 ions and Fe2þ interstitials. Slopes of the linear region >0.5 V were used to calculate defect densities. As can be seen from the polarization scan (Fig. 11(a)) above 0.5 V, the film is not stable and hence this portion of the plot was not used. The defect densities were calculated as per Equation (1) by taking εr as 20. Fig. 13 shows the variation in defect densities for the alloys studied and it can be seen that a decreasing tendency in the defect density is observed with increasing chromium content. Such high defect density concentrations for heavily doped carbon steel were sighted in literature [32]. The indicated high concentration of point defects by MotteSchottky analysis as compared to isolated bulk oxides is attributed to the continuous processes of defect generation and annihilation in the passive films [33]. Similarly, for specimens exposed to controlled DO conditions, MotteSchottky plots showed positive slope, indicating n-type carriers (Fig. 14). Except for 9Cr1Mo, an additional linear region was observed in between 0.0 V and 0.5 V for other alloys. The film was stable in this potential region as indicated by the polarization scan (Fig. 11(b)). In literature, this type of multiple linear regions in
Fig. 12. Mott-Schottky curves for magnetite-coated specimens.
Fig. 13. Variation of defect densities with change in chromium content for specimens exposed to deaerated conditions.
MotteSchottky plots have been assigned to that of deeper donor levels due to metal ions in their higher oxidation state [34,35]. The defect density was calculated by taking εr as 16.5. The defect densities were 2e3 orders less as compared to that of magnetite films. This is attributed to very low Fe2þ in Fe2O3 [36]. The defect density was observed to increase with increase in chromium content in the alloy as given in Fig. 15. For a given alloy, the following table (Table 6) compares the ratio of n-type defect density for oxygenated and deoxygenated conditions. These results showed that the reduction in defect density concentration was more significant in carbon steels (CS and A333) when compared to 2.25Cr-1Mo and 9Cr-1Mo. The smallest change was observed for 9Cr-1Mo. As per the Point Defect Model, defect density in the barrier layer and the migration of the point defects under the influence of the electrostatic field play a major role in the corrosion of structural materials [37]. The metal, oxidized at the metal/oxide interface, is either incorporated in the barrier layer resulting in formation of oxygen vacancies or released as metal ion, into the solution at the oxide/solution interface. The steady state thickness of the barrier layer is achieved when the rate of formation and the dissolution of barrier layer are equal. Corrosion rate under steady state conditions is controlled by cation interstitial injection rate at the metal/oxide interface and barrier layer dissolution rate at the oxide/solution
Fig. 14. Mott-Schottky curves for Hematite-coated specimens.
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deaerated conditions. When the ratios of n-type defect density for oxygenated and deoxygenated conditions were compared for each alloy, it was observed that the reduction in defect density concentration was more significant in carbon steels (CS and A333) than in 2.25Cr-1Mo and 9Cr-1Mo alloys. The minimum change was observed for 9Cr-1Mo. The electrochemical behavior of 9Cr-1Mo under these conditions was attributed to the presence of chromium depletion regions in its microstructure. These studies indicate that Cr-content of the metal and the presence of oxygen greatly influence both the nature of the oxide and its defect chemistry. The oxide properties of A333, a carbon steel containing w0.5% of Cr were found to be superior to that of CS. However further studies are required to understand the influence of flow on the long term corrosion behavior of these alloys. References Fig. 15. Variation of defect densities with change in chromium content for specimens exposed to controlled DO conditions.
interface. In the absence of a steady state, the rate of defect injection (both cation vacancies and oxygen vacancies) at the metal/ oxide interface decides the corrosion rate. In magnetite-type spinel, the presence of Cr decreases the defect density by reducing the rate of injection of vacancies at the metal barrier layer interface, thus reducing the corrosion. Subsequent addition of oxygen in water results in the formation of a hematite layer at the barrier layer/ solution interface. Once this layer is formed, its defect structure/ dissolution rate dominates the corrosion behavior of the underlying metal. The dissolution rate of hematite is several orders of magnitude lower than that of magnetite. This effect of phase change at the barrier layer/solution interface is different from the role played by oxygen at low temperature in increasing the defect density of barrier films on stainless steel [38]. Thus, oxygen addition essentially shifts the location of rate control from the metal/barrier layer interface to the barrier layer/solution interface at the temperature of interest by in situ generation of an inhibitive coating. The decrease in defect densities of carbon steel 106 Gr-B and A333 shows the beneficial effect of Cr content and oxygen addition. 4. Conclusions Preliminary studies were done to assess the effect of chromium on modifying the oxides formed on different chromium containing alloys. For this purpose, four alloys namely CS, A333, 2.25Cr-1Mo and 9Cr-1Mo alloy were chosen. Oxides on these different chromium containing alloys were allowed to grow under the water chemistry conditions similar to PHWR primary heat transport system and their compositions and morphologies were compared with that of carbon steel formed under identical conditions. Chromium gets incorporated in the growing oxide film on chromium containing alloys, in almost the same proportion as in the base material and hence offers more passivity as confirmed by electrochemical studies. The chromium-modified oxide was found to have less defective structure as compared to the oxide on carbon steel. The presence of oxygen aids in hematite formation. This particular film was found to be less soluble with less number of defects as compared to the magnetite film that gets formed under Table 6 Ratio of donor type defects for the specimens. Material
Ratio ND(Fe3O4)/ND(Fe2O3)
CS A333 2.25Cr-1Mo 9Cr-1Mo
w4000 w900 w200 w10
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Contents lists available at ScienceDirect
Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Electrochemical characterization of oxide formed on chromium containing mild steel alloys in LiOH medium Veena Subramanian, Sinu Chandran, H. Subramanian, P. Chandramohan, S. Bera, S. Rangarajan*, S.V. Narasimhan Water and Steam Chemistry Division, Bhabha Atomic Research Centre Facilities, Kalpakkam 603102, Tamilnadu, India
h i g h l i g h t s
g r a p h i c a l a b s t r a c t
High temperature oxides formed on Cr containing mild steels are less defective. Defect densities of oxides decrease with increase in Cr content in the alloy. O2 in solution greatly influences the nature and defect chemistry of oxides.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 June 2013 Received in revised form 10 January 2014 Accepted 1 March 2014
Flow accelerated corrosion leads to wall thinning of outlet-feeder pipes in the primary heat transport system of pressurized heavy water reactors and can even necessitate enmasse feeder replacement. Replacement of carbon steel 106-grade-B (CS) with chromium containing carbon steel reduces the risk of this failure. This paper discusses the role of small additions of chromium in modifying the properties of the oxide film. CS and chromium containing mild steels viz., A333, 2.25Cre1Mo and modified 9Cre1Mo alloy were exposed to primary heat transport (PHT) system chemistry conditions. The oxide films formed were characterized by electrochemical and surface characterization techniques. MotteSchottky analysis showed donor type of defects. The densities of defects in the oxides of chromium containing alloys were 3e15 times less than that in CS. In presence of w200 ppb of dissolved oxygen, the oxides formed were hematite with two orders of magnitude smaller concentration of defects as compared to that formed under reducing conditions. These results suggest that the presence of chromium lowers the defect density of the oxide film and thus ensures a reduced corrosion rate. Ó 2014 Elsevier B.V. All rights reserved.
Keywords: A. Alloys C. Electrochemical techniques C. X-ray photo-emission spectroscopy (XPS) D. Corrosion
1. Introduction In pressurized Heavy Water Reactors (PHWR), the corrosion of the reactor primary circuit materials is minimized by maintaining an alkaline pH (at 25 C) of 10.2e10.6, by the addition of LiOH. Despite maintaining this stringent water chemistry condition, some of the CANDU type of reactors and PHWRs are facing the * Corresponding author. Tel.: þ91 44 27480203; fax: þ91 44 27480097. E-mail addresses: [email protected] (V. Subramanian), [email protected]. in, [email protected] (S. Rangarajan). http://dx.doi.org/10.1016/j.matchemphys.2014.03.003 0254-0584/Ó 2014 Elsevier B.V. All rights reserved.
problem of outlet feeder thinning [1e3]. This phenomenon was attributed to flow accelerated corrosion (FAC). FAC is a process in which carbon steel piping and components corrode in the presence of flowing water or steamewater mixtures with low dissolved oxygen. As water flows over the carbon steel material, stable surface oxide layer (typically Fe3O4) dissolves into the flowing stream, thinning the walls of piping over time and resulting in failures due to rupturing [4e6]. The rate of FAC depends upon hydrodynamic factors such as flow velocity, pipe roughness and geometry of the flow path etc., environmental factors like temperature, pH, iron unsaturation, oxygen concentration and metallurgical factors like
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the composition of steel and residual stress. In addition to the general enhancement in corrosion rate due to high velocity prevailing in feeders, limited area of piping face further enhanced wear rate. This is attributable to pipe geometry which causes localized turbulence and eddies [7]. Hence this type of failure is specific only to outlet feeders of certain PHWRs and CANDUs and to certain channels. Periodic in service inspection campaigns are carried out in reactors to assess the residual life of these piping. However, exhaustive measurements in an operating plant are not always feasible and may even be hampered by access restrictions. So predictive models, which account for fluid dynamics and chemistry at each of the locations, are used to make reliable predictions and aid corrosion assessment of inaccessible areas. Many commercial software like CHECKWORKS, BRT CICERO, RAMEK, COMSY etc. are available to predict FAC [8]. The predictions are based on the assumption that the dominating factor is the presence of magnetite on the corroding surface, its solubility under the prevailing chemistry conditions, mass transfer of the dissolved ions to the fluid and subsequent mass transport to the bulk fluid. These models take in to account the fluid dynamics and chemistry at each of the susceptible locations and make reliable extrapolations for inaccessible areas. The input variables for these models are temperature, pipe geometry, fluid velocity, void fraction, chemical composition of the structural material, pH and dissolved oxygen. Research has shown that trace amounts of alloying elements e particularly chromium, greatly reduce the rate of FAC [9,10]. A study by Cheng et al., demonstrated that an increase in Cr content in the steel can reduce the corrosion rate of steel [11]. The passivity of chromium containing steels is attributed to the formation of a Cr2O3 film. The development of passivity due to incorporation of Cr in MeOeM network of an iron rich alloy requires that the chromium content should be atleast 10% [12e14]. The amount of chromium that is required to be present in the alloy to reduce the FAC rate has been experimentally found to be w0.04% [9,10]. Hence, it is of interest to know how such small amounts of chromium would influence the oxide dissolution and subsequent corrosion behavior. In this paper, an attempt is made to understand the role of chromium in modifying the properties of oxides formed on chromium containing mild steel alloys compared to that of Carbon Steel 106 grade B alloy (CS), which is the current feeder material in some of the Indian PHWRs. ASTM A 333 grade 6 alloy contains about 0.33% of chromium and is proposed as a new material to replace the CS feeders. Hence, this study compares the oxide properties of mild steel alloys viz., A333 (Cr w 0.33%), 2.25Cr1Mo alloy (Cr w 2.25%) and modified 9Cr1Mo (Cr w 9%) along with carbon steel 106 grade B (CS). Oxygen addition to the feedwater is practiced in some pressurized water reactors (PWRs) to protect the carbon steel piping from undergoing FAC, on the secondary side. The added oxygen reacts with the magnetite and converts it to hematite, which has four orders of magnitude lower solubility compared to magnetite. The effectiveness of this modified oxide as a barrier to FAC should also have its origin in the defect density and electronic properties of the oxide film. It is desirable to compare the extent of protection offered by chromium in the alloy and the beneficial effect of added oxygen in the process fluid. This paper also discusses the effect of dissolved oxygen for different chromium containing alloys. The passive films formed at high temperature are semiconducting in nature and the metal release from these oxide films depends on the number density and diffusivity of the ionic point defects [15e18]. Typical donor defects in oxides are oxygen vacancies and cation interstitials and the acceptor defects are mainly cation vacancies. The influence of chromium in changing the defect structure of the oxide was studied by MotteSchottky analysis of the measured capacitance as a function of applied electrode potential.
Table 1 Compositions of the alloys studied.
C Cr Si Ni Mn Mo Cu Co Sn Fe a
Carbon steel
A-333
2.25Cr1Mo
Modified 9Cr1Mo
0.1 <0.3 0.28 <0.14 0.75 <0.03 0.16 <0.03 0.015 Bal
0.24 <0.3 (0.5)a 0.42 <0.14 1.07 e 0.1 <0.013 0.011 Bal
0.065 2.2 <0.2 0.15 0.5 1 0.055 <0.009 <0.008 Bal
0.095 8.1 0.4 0.13 0.38 0.9 0.07 0.025 <0.008 Bal
Measured by atomic absorption spectrometry.
The stability of the film in LiOH and the changes occurring at the oxide/solution interface were studied by electrochemical methods such as Electrochemical Impedance Spectroscopy (EIS) and Potentiodynamic Anodic Polarization (PDAP). 2. Experimental All the chemicals used were GR/AR grade. LiOH was prepared by passing Li2CO3 through OH form of resin. The outlet solution contained 1.5 ppm of Li as analyzed by flame photometry and the pH and conductivity were found to be 10.2 and 45 mS cm1 at 25 C respectively. The three different chromium containing alloys and carbon steel (coupons of approximately 1.2 cm 1.2 cm 0.4 cm) were exposed to lithiated water (LiOH) at 245 C in a static autoclave so as to form their respective oxides for 20 days. The surfaces of the specimens were ground up to 800 grit by using silicon carbide papers, cleaned with ultra-pure water and finally with acetone before exposure. The solution was purged with argon gas for 3e4 h before raising the temperature to ensure removal of dissolved oxygen. One of the parameter that can influence the nature of oxide film is the oxygen content in the solution. Our previous study, wherein the oxide films were formed on these alloys at high temperature by allowing oxygen but only for a short duration, resulted in highly resistive oxides. The composition of the oxides on all the alloys were found to be a mixture of magnetite and hematite in the ratio 3:1 [19]. In the present study, a recirculating autoclave facility was used in order to maintain a steady concentration of oxygen throughout the experiment. In this set of experiments, the specimens were exposed to high temperature simulated PHT conditions, of 245 C and pH25 of 10.2 (by LiOH) in a 2.4 L re-circulating autoclave for 16 days. The flow rate was 10 LPH, thus the water was renewed 4.2 times in an hour. Some of the advantages of switching over to recirculating system from static system were that the solution could always be regenerated and DO could be maintained at 200 10. The chemistry parameters in the source tank, like pH, specific conductivity, dissolved oxygen and redox potential could be measured using an online low temperature chemistry monitoring system. In addition, corrosion potential could be measured online at high temperature and iron-
Table 2 Estimates of oxide thickness on the alloys. Material
Oxide thickness in mm after 20 d exposure to 245 C (deaerated condition)
Oxide thickness in mm after 16 d exposure to 245 C (with DO)
CS A-333 2.25Cr-1Mo Mod. 9Cr1Mo
0.77 0.81 1.0 1.09
1.1 0.91 0.66 0.21
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Fig. 1. SEM micrographs of oxide formed on alloys exposed to deaerated conditions a) CS b) A-333, c) 2.25Cr-1Mo and d) 9Cr-1Mo.
unsaturation could be achieved with the purification flow through a lithiated mixed bed resin at the downstream of the autoclave. Iron-unsaturation in coolant is one of the chemistry parameters which is known to aggravate FAC [20]. The electrochemical potential measured for CS and A333 increased gradually from 0.4 V to 0.1 V and remained steady at 0.1 V after 70 h of exposure. The increase in potential was due to the formation of a stable oxide film. Electrochemical and surface characterization studies of these oxide coated specimens were carried out at room temperature. For
electrochemical studies, the oxide coated specimen (area exposed w1 cm2) was made as the working electrode. The electrolyte used was LiOH, prepared as described above. The electrochemical measurements were made using AUTOLAB PGSTAT 30. Three-electrode configuration was used with a platinum foil serving as the counter electrode and a saturated calomel electrode (SCE) as the reference electrode which was connected through a luggin capillary. In this paper, for the room temperature studies, the values of the potentials are given with respect to SCE only. All the experiments were carried out under deaerated condition by purging the solution with
Fig. 2. SEM micrographs of oxide formed on alloys exposed to oxygenated conditions a) CS b) A 333, c) 2.25Cr-1Mo and d) 9Cr-1Mo.
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Fig. 3. Raman spectra of oxide formed on specimens exposed under deaerated conditions.
argon. A scan rate of 0.5 mV s1 was used for potentiodynamic polarization measurements. Electrochemical impedance spectra were obtained by applying an AC voltage (10 mV amplitude) to the sample at open circuit potential (OCP) in the frequency range of 10 kHze10 mHz. Capacitance measurements on these oxidecoated specimens were carried out from 1.2 V to 0.8 V range and MotteSchottky analyses were performed to assess their defect densities. Scanning electron micrographs (SEM) were taken by using a PHILIPS ESEM XL-30 system. Raman spectra were recorded by using micro Raman Spectrometer (LABRAM HR800) equipped with 514 nm argon ion laser. Laser was focused using 100 objective lens on sample surface with a laser power of 0.5 mW. A grating with 1800 groves/mm was used with a 30 s data acquisition time. X-ray Photoelectron Spectroscopic (XPS) studies were done using VG ESCA Lab MK 200 equipment. The compositions of the alloys were estimated by using direct reading optical emission spectroscopy and are given in Table 1.
3. Results and discussion 3.1. Oxide thickness measurement Estimates of the thickness of the oxide films were made by descaling them using modified Clarke’s solution [21]. Because of its chromium content, oxides on 2.25Cre1Mo and 9Cre1Mo could be descaled only partially using this solution. Hence, a two step process involving caustic permanganate solution (10% NaOH, 4% KMnO4 at 95 C) followed by a treatment with ammonium citrate (10%) was used [22]. The estimates of the thickness of oxides formed on the four alloys, in the presence and absence of dissolved oxygen, are given in Table 2. 3.2. Surface characterization studies 3.2.1. Scanning electron microscopy Scanning electron micrographs of the oxide coated specimens, after exposure to LiOH at 245 C for 20 days under deaerated
Fig. 4. Raman spectra of oxide formed on specimens exposed to controlled oxygenated conditions.
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Fig. 5. XPS spectra of oxide formed at controlled oxygen conditions: (a) Fe 2p and (b) Cr 2p.
conditions, are presented in Fig. 1. The films were made up of crystalline oxide particles and their sizes were found to increase with time of exposure. With all the materials, particularly on the 20 days-exposed specimens, both smaller and bigger oxide crystals were observed, the latter being precipitated from the solution. Energy dispersive X-ray (EDX) microanalysis of these films confirmed that the major constituent was iron. The chromium contents of the films were found to be 0.3%, 0.2%e0.6%, 3%, 6e9% on CS, A-333, 2.25Cre1Mo and 9Cre1Mo alloys respectively. Increase in exposure time did not result in appreciable change in oxide composition for all the materials tested and they are not discussed further in this paper. SEM micrographs of the specimens exposed in the presence of controlled oxygen environment are given in Fig. 2. The surface morphologies of the crystallites were quite different; the typical octahedral edges were not seen as observed on the specimens exposed to deaerated conditions. EDX analysis indicated only the presence of iron and oxygen on the particles. Chromium was seen only on 2.25Cr1Mo and 9Cr1-Mo and the measured concentrations were almost the same as on the base metal. Since the depth of penetration is of the order of w0.5 mm in EDX, the base metal contribution to the chromium content observed in EDX spectra cannot be ruled out.
3.2.2. Raman spectroscopy Raman spectra for the specimens exposed to reducing conditions are given in Fig. 3. The principle peak at 670 cm1 (A1g) is typical of magnetite phase as observed for the oxide formed on CS and A333. On 2.25Cr-1Mo and 9Cr-1Mo, the incorporation of Cr into the ferrite lattice has resulted in a shift of the peak to higher wave numbers. Magnetite is an inverse spinel with all Fe2þ occupying the octahedral (Oh) sites, half of the Fe3þ occupying the tetrahedral (Td) and the remaining Fe3þ occupying Oh sites. It can be represented as [Fe3þ]Td[Fe2þ, Fe3þ]Oh. Here, the Raman active mode is the FeeO stretch involving the [Fe3þ]Td along with the neighboring four oxygen atoms. The strength of the A1g vibrational mode is very sensitive for any substitutions in magnetite. It is reported that the divalent ions like Zn2þ could occupy the Td sites in the magnetite lattice and affect the main peak at w670 cm1 to a greater extent and shift its position significantly [23,24]. However, in Cr3þ substitution, the ion will go and replace the Fe3þ in Oh site, and hence the change brought about in the main Raman active mode will be an indirect one. Hence, the shift could be marginal as observed in the present case. Under controlled oxygenated conditions, all the specimens developed a grayish black oxide film. In their Raman spectra, a
Fig. 6. XPS spectra of oxide formed at deaerated conditions: (a) Fe 2p and (b) Cr 2p.
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distinct and intense peak characteristic of a-Fe2O3 was observed at 227, 248, 294, 412, 500, 613 and 1322 cm1 (Fig. 4). The presence of the oxygen in the water, favored the formation of hematite as an outer layer in all the specimens. In the case of 9Cr1Mo, a small peak at w663 cm1 was seen corresponding to spinel or ferrite/chromite phase.
impedance due to bulk oxide and impedance due to charge transfer processes at oxide/solution interface. Oxide film here can be considered to be made of two distinct layers, the inner film whose composition depends only on the base metal and the outer film which forms due to the precipitation from the solution and hence can vary depending on the additives in the solution. Under the deaerated conditions, the inner and outer oxides are indistinguishable hence
3.2.3. X-ray photoelectron spectroscopy XPS studies on the oxides formed at controlled oxygen conditions confirmed the presence of chromium only on 9Cr-1Mo and iron was present in Fe3þ state on all the specimens (Fig. 5). The Fe 2p3/2 peak was found to be associated with a distinct satellite peak at around 719 eV (marked by an arrow) having a narrow peak width (3.5 eV) which indicated the formation of Fe2O3. For specimens exposed to deaerated conditions, the Fe 2p3/2 peak was broad and the satellite peak characteristic to Fe3þ (around 8 eV left to main peak) was not seen (Fig. 6(a)). In addition, the peak width was around 4.6 eV. All these observations indicated the presence of Fe2þ along with Fe3þ. Under both the conditions, Cr was observed only on 9Cr-1Mo. In both the cases, the Cr 2p3/2 peak was seen at 576.8 eV which indicated the presence of Cr3þ [25]. These observations from Raman and XPS suggested the formation of spinels with Cr. The atomic concentrations of Fe and Cr were evaluated from the peak areas and were found to be 70 at% and 30 at% respectively. Preferential dissolution of Fe, under deareated conditions, probably would have increased the Cr content in the oxide film. Similar observations were made with carbon steel, A333 and 2.25Cr-1Mo specimens. Thus, under deaerated conditions, the surface was dominantly composed of spinels. 3.3. Electrochemical characterization The impedance data was validated using the ‘linear Kramerse Kronig transformation’ function of the commercial software available with the frequency response analysis (FRA) unit from Autolab. Linear KramerseKronig test is based on an approach developed by M.E. Orazem and co-workers and further modified by Boukamp [26]. The matching of the KK transformed and the experimentally obtained data confirmed that the data satisfied all the criteria of a valid impedance data. Electrochemical impedance spectra were measured for all the oxide coated specimens that were exposed to deaerated and oxygenated conditions. The alloys were found to develop a duplex oxide film on their surfaces (Fig. 7). The inner film was found to have smaller particles (w0.1 mm) characteristic of their formation due to the oxidation of the base metal. The outer oxide had relatively bigger particles (1 mm) suggesting their formation due to precipitation from the solution. The outer oxide film could grow with the time of exposure. Hence, the physical model for analyzing EIS data can be represented by a metal covered with a duplex oxide. 3.3.1. EIS for specimens exposed in static autoclave under deaerated conditions EIS data are given in Nyquist and Bode format in Fig. 8(a), (b) and (c) respectively. Analysis of the impedance data was done by fitting them into an electrical equivalent circuit. The total impedance of the system is due to three components. The solution impedance, the
Fig. 7. Schematic of oxide growth as a function of exposure time.
Fig. 8. (a) Nyquist and (b) &(c) Bode plots for oxide covered specimens exposed to deaerated condition.
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505
Fig. 9. Electrical equivalent circuit used for fitting EIS data for specimens exposed to (a) deaerated and (b) oxygenated conditions.
the impedance spectra was found to fit to electrical equivalent circuits with both two and three time constants. Since no significant differences were observed in the c2 value among the fitted data, the simpler circuit with two time constant (Fig. 9(a)) was chosen for this particular case. In Fig. 8, the lines represent the theoretical fit of the experimental data by using this particular equivalent circuit. The impedance response at high frequency corresponds to bulk oxide and at low frequencies to the charge transfer processes occurring at the outer oxide/solution interface. Table 3 gives the fitted impedance parameters. RS is the solution resistance, Rox, the resistance of the oxide film, Cox, the oxide capacitance and Rct, the charge transfer resistance for the processes that may be occurring during the measurement at oxide/solution interface. The Constant Phase Element (CPE), observed is of the order of millifarads and thus corresponds to diffused double layer at the oxide/solution interface. CPE is the non-ideal capacitance and represents a true capacitance if the value of the power term, n, is close to 1. If the exponent value is close 0.5, it indicates diffusion (called Warburg element) as the rate determining process. The value of n away from 1 is also interpreted to represent a nonhomogeneous surface. The bulk oxide properties were observed to be due to the base metal composition as the carbon steel and A333 behaved similarly. Similar trend in bulk oxide properties was observed among the 2.25Cr-1Mo and 9Cr-1Mo alloys also. The oxide capacitance was lower and the resistance higher in 9Cr-1Mo suggesting a larger separation of charges in its oxide. The more resistive interfaces were indicated with increasing Rct values with increasing Cr-content. 3.3.2. EIS for specimens exposed in recirculating loop with controlled dissolved oxygen (DO) The presence of w200 ppb of DO throughout the experiment was ensured in this case. The films formed on all the alloys were found to be hematite. The impedance plots for all the exposed specimens are given in Fig. 10. The inset in the figure gives the impedance spectra at highest frequencies. In this case, the data could not be fitted to the simpler circuit (Fig. 9(a)). In addition to the larger errors it also resulted in the complete loss of information of intermediate frequencies. Hence the equivalent circuit represented in Fig. 9(b) was used. It is quite reasonable since here, the presence of hematite introduces an inhomogeneity in the bulk oxide. It can be seen in Fig. 10 that the experimental data fits quite well with the fitted curve simulating the three time constant equivalent circuit. Table 3 Fitted impedance parameters for specimens exposed to deaerated conditions. Material
Rs (U cm2) Cox (nF cm2) Rox RCT CPE (F cm2) n (U cm2) (kU cm2)
CS 144 A-333 94 2.25Cr-1Mo 257 9Cr-1Mo 301
25.3 24.8 27 14.3
588 604 355 452
59.5 95 234 297
0.3e3 0.2e3 0.3e3 0.6e4
0.8 0.9 0.9 0.8
Table 4 gives the fitted parameters for the impedance data. Rs, Rox, Cox, CPEo/s and Rct have the same meaning as described earlier in Section 3.3.1, for deaerated specimens. In addition, Cpore and Rpore, represents the capacitance and resistance inside the pores at the inner oxide and outer oxide interface. The charge transfer resistances for all the specimens were found to be similar and is fixed at about w500 MU during fitting of the data. The high value suggests a highly resistive oxide/solution interface. For all the specimens, the capacitances were found to be quite similar. The diffused double layer capacitances Cpore and CPEo/s were found to be w2 orders of magnitude lower than CPEo/s at under the deaerataed conditions suggesting a thicker diffused double layer at the interfaces in the oxygenated conditions. This can be attributed to charge distribution in a highly resistive oxide. The oxide film resistances were found to be comparable for all the specimens. The pore resistances of CS and A333 were found be relatively higher than that of 2.25Cr-1Mo and 9Cr-1Mo alloys. A possible reasoning can be given as follows. At lower chromium concentration in the oxide, Cr3þ is likely to be present in the octahedral sites due to the favorable site preference energy [27]. Here, because of the presence of oxygen, the chromium from the outermost layer gets oxidized to chromate and dissolves. Any loss of Cr3þ from the oxide by oxidation to Cr6þ will result in formation of cation vacancies in octahedral sites. Further the charge will be balanced by oxidation of adjacent Fe2þ to Fe3þ. This increased cation vacancy defect density can result in increased transport of the metal ions through the oxide and progressive conversion to a less protective oxide rich in Fe3þ can increase the solubility of the oxide layer, both manifesting as a decreased resistance. Hence the chromium depletion renders the outer oxides on 2.25Cr-1Mo and 9Cr-1Mo alloys more inferior in terms of resistance as compared to CS and A333. 3.3.3. PDAP studies Potentiodynamic anodic scans for specimens previously exposed under deaerated and controlled DO conditions are given in Fig. 11(a) and (b) respectively. The passive current density measured in these scans corresponds to the steady state current and in effect, the steady state concentration of defects in the film. Table 5 gives the range of passive current densities (Ip) deduced from these polarization scans. Ip was observed to decrease as CS > A333 > 2.25Cr-1Mo > 9Cr-1Mo. The Open Circuit Potential (OCP) for hematite coated specimens was observed to be always lower as compared to magnetite coated ones. This shift could be due to the lowering of exchange current density at the cathodic sites at a more resistive oxide. The Ip for hematite films were of the order of nano-amperes as compared to magnetite films which were of the range of micro-amperes. In presence of DO, A333 alloy showed the lowest passive current. This observation essentially confirmed the high resistivity of the films formed with continuous presence of DO in the system.
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Fig. 11. PDAP curves for specimens exposed under (a) deaerated and (b) controlled oxygenated conditions.
3.3.4. Capacitance measurements The magnetite and hematite films formed on the specimens are semiconducting in nature. The metal release from these passive films depends on the concentration of point defects present in them. MotteSchottky analyses were done on the measured capacitance values to calculate defect densities for all the exposed specimens. This method involves the measurement of apparent capacitance as a function of potential under depletion condition where the following relation is obeyed [28,29]:
1 2 kT E E ¼ FB 2 eεr ε0 N e Csc
(1)
where
Fig. 10. (a)Nyquist and (b) and (c) Bode plots for oxide covered specimens exposed to controlled oxygen condition.
Csc ¼ capacitance of the space charge region εr ¼ dielectric constant of the oxide (semiconductor) ε0 ¼ permittivity of free space N ¼ donor density (for n-type)/acceptor density (for p type)
Table 4 Fitted impedance parameters for specimens exposed to oxygenated conditions. Material
Rs (kU cm2)
Cox (nF cm2)
Rox (kU cm2)
Rpore (kU cm2)
Cpore (mF cm2)
RCT (MU cm2)
CPE (F cm2)
n
CS A-333 2.25Cr-1Mo 9Cr-1Mo
0.9 0.9 0.6 0.5
9.9 14.0 11.7 6.4
1.3 1.8 1.1 1.6
2.7 2.8 1.5 1.7
9.1 6.7 10.6 26.7
500 500 500 500
0.2e4 0.3e4 0.3e4 0.5e4
0.8 0.8 0.8 0.8
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Table 5 Ip for the specimens exposed to deaerated and controlled DO conditions. Material
Deaerated
CS A-333 2.25Cr-1Mo 9Cr1-Mo
Controlled DO
OCP (V)
Ip (A cm2) 106
OCP (V)
Ip (A cm2) 109
0.258 0.17 0.212 0.128
1.04 0.47 0.34 0.12
0.402 0.308 0.286 0.14
8000e80 20e0.1 90e3 50e1
E ¼ applied potential EFB ¼ flat band potential e ¼ electronic charge The oxides formed on these alloys were found to follow Motte Schottky relationship at a certain potential range. Fig. 12 shows the typical MotteSchottky plots at 10 Hz for the alloys exposed to deaerated conditions. It was reasonable to take this frequency for measurement, as in our study, at >1 Hz, only the imaginary impedance of the system was found to respond to change in frequency. It is reported in literature that in polycrystalline semiconductor, the MotteSchottky behavior was attributed to the low ionic mobility in the depletion layer, which responds only at low frequencies [30,31]. It can be seen that, except for 9Cr-1Mo, the flat band potential is z0.5 V for all the alloys. Oxides on all the materials show n-type behavior at potentials noble to flat band potential. The donor type defects here are O2 ions and Fe2þ interstitials. Slopes of the linear region >0.5 V were used to calculate defect densities. As can be seen from the polarization scan (Fig. 11(a)) above 0.5 V, the film is not stable and hence this portion of the plot was not used. The defect densities were calculated as per Equation (1) by taking εr as 20. Fig. 13 shows the variation in defect densities for the alloys studied and it can be seen that a decreasing tendency in the defect density is observed with increasing chromium content. Such high defect density concentrations for heavily doped carbon steel were sighted in literature [32]. The indicated high concentration of point defects by MotteSchottky analysis as compared to isolated bulk oxides is attributed to the continuous processes of defect generation and annihilation in the passive films [33]. Similarly, for specimens exposed to controlled DO conditions, MotteSchottky plots showed positive slope, indicating n-type carriers (Fig. 14). Except for 9Cr1Mo, an additional linear region was observed in between 0.0 V and 0.5 V for other alloys. The film was stable in this potential region as indicated by the polarization scan (Fig. 11(b)). In literature, this type of multiple linear regions in
Fig. 12. Mott-Schottky curves for magnetite-coated specimens.
Fig. 13. Variation of defect densities with change in chromium content for specimens exposed to deaerated conditions.
MotteSchottky plots have been assigned to that of deeper donor levels due to metal ions in their higher oxidation state [34,35]. The defect density was calculated by taking εr as 16.5. The defect densities were 2e3 orders less as compared to that of magnetite films. This is attributed to very low Fe2þ in Fe2O3 [36]. The defect density was observed to increase with increase in chromium content in the alloy as given in Fig. 15. For a given alloy, the following table (Table 6) compares the ratio of n-type defect density for oxygenated and deoxygenated conditions. These results showed that the reduction in defect density concentration was more significant in carbon steels (CS and A333) when compared to 2.25Cr-1Mo and 9Cr-1Mo. The smallest change was observed for 9Cr-1Mo. As per the Point Defect Model, defect density in the barrier layer and the migration of the point defects under the influence of the electrostatic field play a major role in the corrosion of structural materials [37]. The metal, oxidized at the metal/oxide interface, is either incorporated in the barrier layer resulting in formation of oxygen vacancies or released as metal ion, into the solution at the oxide/solution interface. The steady state thickness of the barrier layer is achieved when the rate of formation and the dissolution of barrier layer are equal. Corrosion rate under steady state conditions is controlled by cation interstitial injection rate at the metal/oxide interface and barrier layer dissolution rate at the oxide/solution
Fig. 14. Mott-Schottky curves for Hematite-coated specimens.
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deaerated conditions. When the ratios of n-type defect density for oxygenated and deoxygenated conditions were compared for each alloy, it was observed that the reduction in defect density concentration was more significant in carbon steels (CS and A333) than in 2.25Cr-1Mo and 9Cr-1Mo alloys. The minimum change was observed for 9Cr-1Mo. The electrochemical behavior of 9Cr-1Mo under these conditions was attributed to the presence of chromium depletion regions in its microstructure. These studies indicate that Cr-content of the metal and the presence of oxygen greatly influence both the nature of the oxide and its defect chemistry. The oxide properties of A333, a carbon steel containing w0.5% of Cr were found to be superior to that of CS. However further studies are required to understand the influence of flow on the long term corrosion behavior of these alloys. References Fig. 15. Variation of defect densities with change in chromium content for specimens exposed to controlled DO conditions.
interface. In the absence of a steady state, the rate of defect injection (both cation vacancies and oxygen vacancies) at the metal/ oxide interface decides the corrosion rate. In magnetite-type spinel, the presence of Cr decreases the defect density by reducing the rate of injection of vacancies at the metal barrier layer interface, thus reducing the corrosion. Subsequent addition of oxygen in water results in the formation of a hematite layer at the barrier layer/ solution interface. Once this layer is formed, its defect structure/ dissolution rate dominates the corrosion behavior of the underlying metal. The dissolution rate of hematite is several orders of magnitude lower than that of magnetite. This effect of phase change at the barrier layer/solution interface is different from the role played by oxygen at low temperature in increasing the defect density of barrier films on stainless steel [38]. Thus, oxygen addition essentially shifts the location of rate control from the metal/barrier layer interface to the barrier layer/solution interface at the temperature of interest by in situ generation of an inhibitive coating. The decrease in defect densities of carbon steel 106 Gr-B and A333 shows the beneficial effect of Cr content and oxygen addition. 4. Conclusions Preliminary studies were done to assess the effect of chromium on modifying the oxides formed on different chromium containing alloys. For this purpose, four alloys namely CS, A333, 2.25Cr-1Mo and 9Cr-1Mo alloy were chosen. Oxides on these different chromium containing alloys were allowed to grow under the water chemistry conditions similar to PHWR primary heat transport system and their compositions and morphologies were compared with that of carbon steel formed under identical conditions. Chromium gets incorporated in the growing oxide film on chromium containing alloys, in almost the same proportion as in the base material and hence offers more passivity as confirmed by electrochemical studies. The chromium-modified oxide was found to have less defective structure as compared to the oxide on carbon steel. The presence of oxygen aids in hematite formation. This particular film was found to be less soluble with less number of defects as compared to the magnetite film that gets formed under Table 6 Ratio of donor type defects for the specimens. Material
Ratio ND(Fe3O4)/ND(Fe2O3)
CS A333 2.25Cr-1Mo 9Cr-1Mo
w4000 w900 w200 w10
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