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Corrosion of steel alloys in eutectic NaCl+Na2CO3 at 700 °C and Li2CO3 + K2CO3 + Na2CO3 at 450 °C for thermal energy storage

Madjid Sarvghad*, Theodore A. Steinberg, Geoffrey Will

Science and Engineering Faculty, Queensland University of Technology (QUT), Queensland, Australia 4001

Abstract

Stainless steel 316, duplex steel 2205 and carbon steel 1008 were examined for compatibility with the eutectic mixtures of NaCl+Na2CO3 at 700 °C and Li2CO3 + K2CO3 + Na2CO3 at 450 °C in air for thermal energy storage. Electrochemical measurements combined with advanced microscopy and microanalysis techniques were employed. Oxidation was found as the primary attack in both molten salt environments. However, the availability of O2 controlled the degree of oxidation. Alloy 316 showed the lowest corrosion current densities in each molten salt owing to the formation of films on the surface. The attack morphology on the surface of all materials was uniform corrosion with no localized degradation at 450 ºC. Microscopy observations showed grain boundary oxidative attack as the primary corrosion mechanism for all studied alloys at 700 °C with depletion of alloying elements from grain boundaries in alloys 316 and 2205. The protective nature of austenite phase reduced selective oxidation of the underlying ferrite layers of alloy 2205 in chloride carbonate at 700 °C.

Keywords

Steel; Molten Salt; Corrosion; Impedance spectroscopy; Polarization; Microscopy

                                                            * Corresponding author; Email: Sarvghad_madjid@hotmail.com; madjid.sarvghadmoghaddam@hdr.qut.edu.au  

 

 

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Graphical abstract

 

 

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1. Introduction

Recent interest in increasing plant efficiency requires improved compatibility between structural alloys and molten salts for Thermal Energy Storage (TES) used in Concentrated Solar Thermal (CST) power plants [1-3]. Hot corrosion and oxidation from retained storage media (molten salts) as Phase Change Materials (PCM) causes significant deterioration to containment materials used in TES systems [1-14]. Whilst steel alloys are considered as economic candidates for containment materials in TES systems, eutectic compositions of molten chlorides, nitrates, carbonates and fluorides are favorite candidates as storage medium [15-20]. Oxyanions like carbonates have recently attracted more attention not only owing to their great heat capacity, but also for their high energy density and less corrosive behavior compared to traditional chlorides [2, 21]. Even the mixture of carbonate and chloride have been reported to be less corrosive than pure chlorides [21].

The compatibility of structural alloys with molten salts depends on the metal’s potential for oxidation, passivating nature and the solubility of corrosion products in the salt [22]. In steel alloys, scaling rate depends on three variables: steel chemistry, temperature and atmosphere [23]. Based on the working temperature low-alloyed carbon steels, Cr-Mo steels and Cr-Ni stainless steels might be employed in combination with molten salts [10]. Groll et al. [24] reported intergranular corrosion of steel alloys in contact with molten chloride salts at temperatures up to 420 °C. Other studies confirm that intergranular attack in Fe-Ni-Cr alloys is more severe than metal loss in molten chlorides [21]. Ferritic steel has shown to oxidize more easily than austenitic steel in hot chloride containing atmospheres [25]. However, recent research on a group of metal alloys also showed acceptable resistance of stainless steel 310 to molten carbonate salts at 750 °C [2]. Steel alloys with around 20 wt% Cr and/or high nickel content show a greater resistance to high-temperature corrosion [23, 26].

The stability of oxide layer should also be taken into consideration when using carbon or stainless steels. Most metals oxidize over a wide range of conditions at elevated temperatures. Thus, oxidation rate and morphology are of importance to determine the material lifetime. Corrosion resistance in many high-temperature environments is achieved by the formation of a protective oxide film on the material surface [27]. An early study by Azzi et al. [28] showed that the corrosion rate of iron in molten carbonate is limited by the diffusion of oxidizing species through the corrosion products rather than in the melt; where FeO, Fe2O3 and Fe3O4 were determined as the general corrosion products formed on the material surface. Hot corrosion behavior of steel alloys in molten Li2CO3-K2CO3 showed the formation of a porous layer composed of Fe2O3 and LiFe5O8 on the material surface [29]. However, the usually thick Fe2O3 layer has shown not to be protective in molten salt environments [30]. Al and Si also contribute to the development of self-healing protective oxide films on the alloy surface acting as diffusion barriers against further oxidation [15]. Cr and Al oxides have also proved to be more stable than those of iron and thus Cr2O3 and Al2O3 are thermodynamically favored [31].

The compatibility of containment material with molten salt and its stability is a concerning issue in TES systems [8]. The selection of appropriate and optimum structural materials as Thermal Energy Storage vessel, subject to corrosive molten salts as PCM and high temperatures atmospheres, is essential in developing economic and functionally efficient systems. This study will examine the corrosion behavior of three commercial steel alloys in two eutectic mixtures of molten salts for the next generation of TES applications.

2. Experimental procedure

Austenitic stainless steel 316 (SS316), ferritic carbon steel 1008 (CS1008) and ferritic/austenitic duplex steel 2205 (DS2205) were examined in two eutectic mixtures of molten salts. Table 1 summarizes the structures and nominal compositions of the alloys used in the current study.

 

 

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Sodium carbonate anhydrous LR (CAS No. 497-19-8), sodium chloride AR (CAS No. 7647-14-5), potassium carbonate anhydrous LR (CAS No. 584-08-7) and lithium carbonate 99% (CAS No. 554-13-2) were placed for 24 h in a 180 °C furnace to dry and then were measured and mixed according to Table 2. The eutectic mixtures melting points and test temperatures are also provided in the table. Test temperature for each salt was selected close to its melting point based on the assumption that the salt will be used as a PCM in a CST plant.

Table 1 Nominal elemental composition and crystal structure of the studied alloys (wt%).

Table 2 Chemical composition, melting point and test temperature of salt mixtures.

Salt mixture Composition (wt%) Melting point

(°C) [32] Test temperature

(°C) Chloride carbonate 40 NaCl + 60 Na2CO3 632 700 Ternary carbonate 33.4 Na2CO3 + 32.1 Li2CO3+ 34.5 K2CO3 397 450

2.1. Electrochemical corrosion investigation

Electrochemical experiments were conducted using a three-electrode cell containing the molten salts in alumina crucibles open to air at 700 °C and 450 °C in a preheated cylindrical furnace. Test coupons of 25 mm long, 5 mm wide and 1.4 mm thick were mechanically wet ground and polished down to 0.04 m by colloidal silica, washed with ethanol and dried in air. Measurements were implemented by means of a VMP3-based BioLogic instrument controlled by EC-Lab® software. The three-electrode cell was implemented with the polished sample as the working electrode and two same sized platinum sheets (25×5×1 mm) as pseudo reference and counter electrodes [33-36]. Samples were subjected to open circuit potential (OCP), electrochemical impedance spectroscopy (EIS) and potentiodynamic polarization (PDP) measurements in this order to avoid sample deterioration. Equilibrations of potentials (OCP) were carried out for 1 h immediately after immersion. EIS measurements were then obtained using a frequency range of 100 kHz- 100 mHz with the amplitude of ±10 mV. Finally, PDP was conducted at the potential scan rate of 10 mV/min and potential range of -400 to +500 mV with respect to the open circuit potential. ZFit analysis of EC-Lab software was used to fit successive impedance cycles.

2.2. Static corrosion

Fresh metal coupons were cut to around 25×7×1.4 mm for static corrosion tests while the front sides were mechanically wet polished down to 1 µm in colloidal silica using standard grinding and polishing procedures. A schematic representation of the test condition and alignment of samples in the furnace is shown in Fig. 1. Cylindrical alumina crucibles were used as salt vessels and the furnace temperature was set to 700 ± 10 °C for chloride carbonate and 450 ± 10 °C for ternary carbonate.

The salt containing vessels were placed into the furnace at room temperature and then gradually heated up to the test temperature. Once the salt melted and the chamber conditions stabilized, the metal coupons were immersed so that the top half was exposed to the air and bottom half submerged into the molten salt. Such a configuration enabled us to make a comparison not only between the impact of the salt and that of the oxygen from the environment, but also the decomposition gases of the salt which are expected to increase the corrosive impact in this atmosphere. Metal coupons were removed after 120 h of exposure. All coupons were then mounted exposing the side indicated in Fig. 1 into a conductive resin,

Alloy Structure Fe Ni Cr C Mo Mn S P SS316 fcc Bal 12 17.5 0.05 3 … … …

DS2205 fcc/bcc Bal 6.5 23 0.03 3.5 … … … CS1008 bcc Bal … … 0.14 … 0.5 0.04 0.04

 

 

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ground and polished down to 0.04 µm from the side of the sample in colloidal silica using standard procedures, washed with ethanol and finally dried in air. Microscopy points in Fig. 1 provide us with a view of the material close to the polished side. It is worthy to note that although 120 h does not seem long enough to study corrosion rate and related phenomena, aggressive nature of molten salts makes studying short-term impacts like attack morphology and corrosion mechanisms possible.

Fig. 1 Schematic of the corrosion vessel and a sample inside the furnace.

2.3. Metallography samples

Small coupons of each alloy were selected and mechanically wet ground and polished down to 0.04 μm by colloidal silica, washed with ethanol and finally dried in air. These coupons were used to study the corroded metal morphology and short-term impacts of the molten salts on the microstructure. The coupons were submerged into the molten salts for only 3 min and then were ultrasonically cleaned for 15 min in demineralized water for corrosion product residues to be removed from the surface. The samples were then studied under an optic microscope from the surface.

2.4. Macro and micro-structural investigations 2.4.1. Optical microscopy analysis

Optic microscopes model Leica DMi8A (magnification 1.25x-50x) and Leica M125 (magnification 0.8x-10x) both equipped with Leica Application Suite software were used to take macro and micro-images for microstructural and corrosion observations.

2.4.2. SEM, EDS and EBSD analysis

Complimentary techniques like scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS) and electron back-scatter diffraction (EBSD) were employed for further microstructural investigations using a field emission SEM (model: JEOL 7001F, with automated feature detection equipped with secondary electron, EDS analysis system, OXFORD EBSD pattern analyzer and Channel 5 analysis software).

3. Results and discussion 3.1. Electrochemistry

OCP graphs of the alloy specimens in the studied molten salts are represented in Fig. 2a. CS1008 shows stable OCP values over the 1 h of exposure in both salts after an initial rapid drop in ternary carbonate. DS2205 and SS316 show more noble potential values over time in chloride carbonate which could be

 

 

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attributed to the gradual development of a film on the surface. These alloys also show steady OCP values in ternary carbonate.

Fig. 2 (a) Open circuit potential, and (b) potentiodynamic polarization curves of samples in molten chloride carbonate at 700 °C and ternary carbonate at 450 °C.

Ecorr values in PDP plots in Fig. 2b reflect OCP values in Fig. 2a. Data extracted from PDP curves are summarized in Tables 3 and 4 depicting the corrosion current density values in Fig. 3. SS316 represents the best performance and shows the lowest current densities in each salt. Graphs show that DS2205 and SS316 perform similarly in each salt with similar potential and comparable Icorr values. Higher Icorr value of DS2205 compared to that of SS316 in ternary carbonate could be attributed to the faster rates of cathodic and anodic reactions according to Fig. 2b and Table 4. On the other hand, maximum Icorr values belong to CS1008 in both environments. Ecorr for CS1008 is comparable however, the significant change in Icorr could indicate a change in mechanism between the two molten salts.

Table 3 Extracted data from polarization plots in Fig. 2b for the alloys in chloride carbonate at 700 ºC. Alloy Ecorr (mV)

Icorr (µA.cm-²)

Βa (mV/decade)

|Βc| (mV/decade)

SS316 -24.0 247.2 220.9 312.9DS2205 -35.6 293.5 239.4 360.6 CS1008 -920.5 1686.5 163.1 409.0

Table 4 Extracted data from polarization plots in Fig. 2b for the alloys in ternary carbonate at 450 ºC. Alloy Ecorr (mV)

Icorr (µA.cm-²)

Βa (mV/decade)

|Βc| (mV/decade)

SS316 -538.5 5.7 55.9 142.6 DS2205 -547.1 29.7 100.3 299.8CS1008 -902.5 85.2 153.6 264.6

 

 

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Fig. 3 Column charts of corrosion current density values in chloride carbonate at 700 °C (a) and ternary carbonate at 450 °C (b).

It is known from literature that molten carbonate corrosion becomes less severe at lower temperatures and the corrosion mechanism can vary widely based on temperature and salt composition [21]. Regardless of the salt composition and test temperature, early studies on the molten salt corrosion of metals express the general corrosion reaction as a combination of metal dissolution, partial anodic reaction and reduction of oxidants, and partial cathodic reaction [37]. Liquid salts are ionic and their interaction with metals is therefore electrochemical [38-40]. Consequently, the partial anodic reaction could be considered as the oxidation of the metal [37, 39]:

M→Mn ne- Equation 1 2Mn nO2-→2MOn Equation 2 where M is the transition metal like Fe, Cr and Ni. Partial cathodic reaction follows the equation below [37]:

Ox ne-→R Equation 3 with n pointing to the number of electrons while Ox and R refer to oxidant and reductant, respectively.

Corrosion mechanisms of metals in molten carbonate salts have been earlier discussed in detail [41]. It has been also reported that the solubility of a range of oxides in molten salts highly depends on the melt basicity and so does the type of dissolution mechanism [29]. Therefore, corrosion will be dependent upon the solubility of the formed metal oxides on the material surface. Azzi et al. [28] reported that oxygen can be dissolved in molten Na2CO3-K2CO3 at 800 °C. Fe2+ ions were then reported to be formed in the melt due to the dissolution of oxidizing species and the simultaneous oxidizing of iron according to reaction 1. The presence of oxygen causes the formation of peroxide ions in the melt leading to higher corrosion rate. Oxygen participates indirectly in the corrosion of iron by O peroxide formed via equilibrium 4 and reduced to oxygen anions according to reaction 5 [28]:

12 O2 O2-↔O22- Equation 4 O22- 2e-→2O2- Equation 5 Early studies [37, 42, 43] suggest that highly reactive peroxide ions O and superoxide ions O2- could be considered as the basic component of the solvent in molten salts, which oxidize the metal when reduced on the metal surface. In O2 atmosphere, reaction 5 participates in the corrosion process and

 

 

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leads to the formation of iron oxide products like FeO, Fe2O3 and Fe3O4 [28, 29]. Therefore, high concentration of oxide ions in the melt leads to the assumption that the carbonate molten salts used in this study act as a basic melt implying that the corrosion of the metal is under anodic control [37].

Kruizenga et al. [44], on the other hand, reported that stable passivated oxide scales are not likely to form on structural alloys in molten chloride salts. It is already known from literature that, thermodynamically, alkali chlorides are more favored than transition metal chlorides [15]. Consequently, as the common alloying elements are not expected to reduce the salt, corrosion is expected to be driven by oxidants like H2O, O2, H+ and other impurities in molten chlorides [15, 37, 44]. Therefore, any moisture content in the system could lead to the formation of HCl and chlorine gas [15, 40, 45].

Despite of the presence of HCl in the system, Cl2 has been found as the main corrosive medium because of its ability to penetrate protective oxide scales and react with transition metals [40, 46]. Therefore, the production of chlorine gas and its reaction with metal oxides on the surface could be defined as the dominant corrosion mechanism in molten chlorides [40].

For further analyses, EIS was measured in the molten salts; Nyquist and Bode-Phase plots are shown in Fig.4. Simulated equivalent circuit (EC) is included in Fig. 4a and the data extracted from the model is summarized in Table 5. Here RS, Rct and Rox point to solution resistance, charge transfer (polarization) resistance and transfer resistance of ions in the oxide film, respectively. Qdl is the double layer Constant Phase Element (CPE), Qox is the film CPE, n represents the CPE parameter and W corresponds to Warburg resistance. As reported previously [41], CPE points to the non-ideal capacitance of the surface and is used to calculate the surface charge transfer capacitance during corrosion [47-49].

In chloride carbonate at 700 ºC, the presence of CPE (Fig. 4a) confirms the formation of scale layers and subsequent microscopic roughness on the metal’s surfaces where the transfer of ions in the scale is rate limiting [50, 51]. This is also in agreement with OCP plots in Fig. 2a. However, the appearance of Warburg mass transfer impedance indicates corrosion due to the non protective film on the surface [51]. Warburg coefficient is inversely proportional to the square root of diffusion coefficient [52]. Therefore, the highest W coefficient for DS2205 means lowest diffusion through the surface scale for this alloy compared to that of the SS316 with the highest diffusion coefficient. Table 5 and Bode plots in Fig. 4b show the highest impedance values for DS2205 at low frequencies. Therefore, probably due to the high Cr content in the alloy composition, a semi-protective film with very low diffusivity on the surface of DS2205 should have been developed [29, 53]. In addition, as will be shown later, alloy CS1008 shows the development of a very thick oxide layer on the surface which does not provide protection against further corrosion. According to Table 5, the diffusivity of this layer should be quite low.

In ternary carbonate at 450 ºC, the EC model confirms the formation of scale layers on the surface of the alloys (Fig. 4c). Again, the appearance of Warburg mass transfer impedance casts a shadow of doubt over the protective nature of the film and points to high corrosion rates for all studied alloys. Once more, DS2205 shows the highest impedance values, according to Fig. 4d and Table 5, which could be a result of the formation of a semi-protective film on the surface due to high Cr content in the alloy. In agreement with its highest corrosion current density value, CS1008 shows the lowest impedance values while SS316 displays moderate values. Contrary to the chloride carbonate salt, W coefficient values here point to the lowest diffusivity through the surface layer on SS316 versus the highest in DS2205. However, these values are comparable in this ternary carbonate salt.

 

 

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Fig. 4 (a) Nyquist, and (b) Bode-Phase plots resulted from impedance measurements of the alloys in chloride carbonate at 700 °C, (c) Nyquist, and (d) Bode-Phase plots resulted from impedance measurements of the alloys in ternary carbonate at 450 °C. Larger Nyquist plots of CS1008, calculated equivalent circuit and fit data points are also included in (a) and (c).

Table 5 EIS data extracted from equivalent circuit models in Fig. 4.

3.2. Microscopy and microanalysis

Optical and SEM images of samples submerged for 3 min in molten chloride carbonate at 700 ºC are presented in Fig. 5. An adhesive layer was formed on DS2205 and CS1008 while the deposit on SS316 was physically detached during handling and washing. The film on CS1008 could be removed after 15 min ultrasonic cleaning with the film on DS2205 remaining attached to the surface, Figs. 5c and f.

Molten Salt Alloy

Rs Rct Rox Qdl ndl

Qox nox

W (Ω.cm2) (Ω.cm2) (Ω.cm2) (F.sn-1.cm-2) (F.sn-1.cm-2) (Ω.s-1/2)

Chloride carbonate

SS316 1.03 1.60 9.72 0.19 0.47 0.05 1.00 8.96

DS2205 0.62 32.1 40.1 0.05 0.64 0.05 0.68 1.1e+6

CS1008 0.97 0.32 10.95 0.11 0.57 0.43 0.49 6.9e+5

Ternary carbonate

SS316 1.61 3.10 204.38 1.24e-03 0.76 3.12e-03 0.79 69.01

DS2205 1.23 543.85 5.76 3.68e-03 0.77 8.15e-04 0.86 31.02

CS1008 1.21 3.08 2.86 3.50e-03 0.50 5.78e-03 0.30 44.27

 

 

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Grain boundary (GB) attack seems to be the dominant corrosion mechanism for all studied alloys at 700 °C. Even on DS2205, surface contamination seems to follow grain boundaries, Fig. 5f. This is more obvious in the SEM image in Fig. 5e where salt residues follow GBs on alloy CS1008. However, as will be shown later, the oxidation is so severe on CS1008 that after long exposure times the morphology looks like uniform rather than GB attack. Further detailed microstructural observations after long-term exposure to each eutectic salt are provided later.

Fig. 5 Optical (a to c) and SEM (d to f) images of as-polished samples after 3 min exposure to the molten chloride carbonate at 700 °C. Inserted EBSD phase color map of as-received material and parallel features in (c) point to the alternate layers of ferrite/austenite in the texture of the duplex steel 2205 corroded with different rates.

3.2.1. SS316

An optical image of a coupon of SS316 subjected to chloride carbonate at 700 °C for 120 h is shown in Fig. 6a. Hanging the sample over the salt bath resulted in the top half being exposed to air, salt vapor and potentially salt creep and the bottom half to the molten salt. Figs. 6b and c compare and contrast top and bottom halves of the sample, respectively. Images confirm GB attack in both environments. In the absence of a reference surface and ignoring the material already removed through corrosion, attack above the salt penetrates around 54.4 µm towards the bulk material (Fig. 6b) versus 35.1 µm at the bottom part under the salt level (Fig. 6c). No adherent film is detectable on the top part while a semi-protective deposit seems to be formed under the salt level. However, no film could be observed at the bottom side of the sample under the salt where the penetration depth of 57.5 µm is close to that of the top part exposed to air. Hence, lower GB degradation depth under the salt level could be attributed to the formation of a semi-protective film on the surface as previously predicted by electrochemistry.

 

 

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Fig. 6 (a) Macro-graph of a SS316 sample hung over the molten chloride carbonate at 700 °C for 120 h, (b) optical image of the alloy cross-section from sidewall above the molten salt, (c) optical image of the alloy cross-section from sidewall under the salt level.

Further SEM-EDS analyses in Figs. 7 and 8 clarify the behavior. 54.4 µm penetration depth above the molten salt accompanies Fe and Cr depletion from GBs and the following GB oxidation. The lower penetration depth under the salt level also follows the formation of GB oxides, Fig. 8. Therefore, GB degradation due to the de-alloying of Fe and Cr from GBs and formation of Fe-Cr oxides could be considered as the main corrosion mechanism which, in the presence of the molten salt, is hindered owing to the development of a semi-protective film on the surface. This is in agreement with electrochemical measurements reported previously. The reduced amount of oxygen in the salt, compared to the above atmosphere, could have also contributed. However, as previously shown in Figs. 5a and d and according to the high diffusivity previously observed through the surface layer, the oxide film does not seem adherent as it might have been detached from the bottom side in the salt and also from the top side exposed to air due to gravity or during handling.

 

 

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Fig. 7 (a) SEM image, and (b) to (f) its corresponding EDS map analysis of an area on the top corner of SS316 above the chloride carbonate salt level and exposed to air for 120 h at 700 °C.

Fig. 8 (a) SEM image, and (b) to (f) its corresponding EDS map analysis of an area on the bottom corner of SS316 submerged into the chloride carbonate salt for 120 h at 700 °C.

Fig. 9 shows the formation of a continuous and adherent Fe-Ni-Cr-K oxide on the material surface in ternary carbonate at 450 °C. No localized de-alloying and oxidation/corrosion attack is detectable on the sample without any noticeable discrepancy between the attack depth above and below the salt level. Even the possible formation of lithium oxide, according to the fact that Li is hard to be detected by EDS [54], does not seem to have contributed to corrosion under the salt level. Thus, it could be concluded that the oxide layer formed on the material surface acts as a solid barrier against further corrosion in contact with the molten salt at 450 °C.

 

 

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Fig. 9 (a) Optical image of the top corner of SS316 above the ternary carbonate salt and exposed to air at 450 °C for 120 h, (b) SEM image of the red rectangular area in (a), (c) its corresponding EDS map analysis, (d) optical image of the bottom corner of SS316 under the ternary salt at 450 °C for 120 h, (e) SEM image of the red rectangular area in (d), and (f) its corresponding EDS map analysis.

3.2.2. CS1008

Degradation of the ferritic carbon steel CS1008 above the chloride carbonate salt at 700 °C for 120 h accompanies the formation of a thick layer of iron oxide on the surface, Figs. 10a and b. The deposit seems to consist of two separate oxide layers with the total thickness of 1.16 mm on the sample surface which was exposed above the salt level. Although a similar oxide seems to have been developed on the material surface under the salt level, its total thickness of 0.80 mm is 31% less than that of the above; Fig. 10c. On the other hand, 0.80 mm thickness of the metal remained intact after 120 h exposure is 21% higher than that above the salt (0.63 mm). EDS analysis confirms the penetration of chlorine through the film and subsequent chlorine attack as discussed previously. Therefore, oxidation seems as the main corrosion behavior for CS1008 in contact with chloride carbonate salt at 700 °C while the molten salt here seems to have played a protective role against further oxidation probably due to the lower O2 in the salt.

Corrosion of CS1008 in ternary carbonate at 450 °C for 120 h is shown in Fig. 11. According to Fig. 11a, a 50.5 µm thick and adherent film of iron oxide has been formed on the alloy surface for the portion of the sample above the molten salt. The bottom half of the metal submerged into the molten salt shows an adherent film which appears thinner; Figs. 11c and d. EDS analysis shows the similar film composition of iron oxide formed under the salt compared to the above. Therefore, oxidation could be defined as the main corrosion attack in the ternary carbonate at 450 °C with the salt decelerating the development of the oxide layer.

 

 

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Fig. 10 (a) SEM image of an area on the top half of CS1008 above the chloride carbonate salt exposed to air for 120 h at 700 °C, (b) SEM-EDS map analysis corresponding to the red rectangular area in (a), (c) SEM image of an area on the bottom half of CS1008 under the chloride carbonate salt for 120 h at 700 °C, (d) SEM-EDS map analysis corresponding to the red rectangular area in (c).

Fig. 11 (a) Optical image of an area on the top corner of CS1008 above the ternary carbonate salt at 450 °C for 120 h, (b) SEM-EDS map analysis corresponding to the red rectangular area in (a), (c) optical image of an area on the bottom corner of CS1008 submerged into the ternary carbonate salt at 450 °C for 120 h, (d) SEM-EDS map analysis corresponding to the red rectangular area in (c).

 

 

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3.2.3. DS2205

Although the corrosion current density values of alloy DS2205 is closer to that of the SS316, this dual phase ferritic-austenitic steel inherits the combined behavior of SS316 and CS1008 regarding the corroded metal morphology in the studied molten salt environments. Fig. 12 shows oxidation attack of ferrite above the chloride carbonate salt at 700 °C for 120 h. Despite the outcomes of electrochemical measurements, no surface deposit is detectable on the alloy surface above and under the salt level, Figs. 12a and d. In fact, as shown in Figs. 5c and f, a surface film is likely to form on the alloy surface in contact with the salt. However, the film could have been later dissolved as a result of the its solubility in the salt (or its vapors) due to the acidic nature of the molten chloride as discussed earlier [40].

EBSD phase and elemental analysis in Fig. 12c show selective oxidation of ferrite, rich in Ni-Cr-Mo, while the alternate austenite phase remained relatively intact. Optical and EBSD images in Figs. 12d to f show the same behavior on the sample under the salt level.

Fig. 12 (a) Optical image from top corner of DS2205 above the chloride carbonate salt and exposed to air for 120 h at 700 °C, (b) EBSD band contrast image of the red rectangular area in (a), (c) EBSD phase and elemental analysis maps corresponding to (b), (d) optical image from bottom corner of DS2205 under the chloride carbonate salt for 120 h at 700 °C, (e) EBSD band contrast image of the red rectangular area in (d), (f) EBSD phase and elemental analysis maps corresponding to (e).

Again, ignoring the material already removed through corrosion, the oxidation depth from the top corner above the salt is around 23.2 µm in the direction normal to alternate austenite/ferrite phases (rolling direction) while it extends roughly twice that distance (up to 43.1 µm) on the surface parallel to the rolling direction; Figs. 13a and b. In fact, oxidation proceeds faster along ferrite while austenite acts as a barrier against further oxidation and protects the underlying ferrite layer. However, the presence of austenite as a more noble phase at the vicinity of ferrite (ref. Fig. 2a) might have contributed to faster oxidation due to the galvanic corrosion of ferrite which is less noble than austenite [25, 41, 55, 56]. This is similar to that part of the metal under the salt level, Figs. 13c and d. The 20.5 µm penetration depth from the edge is only 11.6% lower than that of the above. However, compared to above the salt level (Fig. 13b), it seems that the limited solubility of oxygen in the salt has reduced the oxidation by more than 50% parallel to the rolling direction below the salt.

 

 

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Fig. 13 (a) and (b) SEM images from the side edge and top edge, respectively, on the top part of DS2205 above the chloride carbonate salt exposed to air for 120 h at 700 °C, (c) and (d) SEM images from the side edge and bottom edge, respectively, on the bottom part of DS2205 under the chloride carbonate salt for 120 h at 700 °C.

Fig. 14 shows the top part of alloy 2205 above the ternary carbonate salt after 120 h at 450 °C suffered from oxidation on the surface. The degradation depth is around 22 µm. On the other hand and having in mind that the sample surface was polished before exposure, a sharp edge is detectable in Fig. 15 for the front side of the sample in direct contact with the molten salt. That sharpness indicates that the metal remained intact under the salt. Comparing Fig. 15 to the previous observations, it seems that 120 h was not long enough for considerable corrosion attack in the ternary carbonate salt at 450 °C. This also supports the salt’s protection role because of the reduced availability of O2 under the salt.

Fig. 14 (a) SEM image from the top part of DS2205 above the ternary carbonate salt exposed to air at 450 °C for 120 h, (b) EDS map analysis corresponding to (a).

 

 

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Fig. 15 SEM image from the bottom part of DS2205 below the ternary carbonate salt at 450 °C for 120 h.

Conclusion

Corrosion behavior of three commercial alloys including stainless steel 316, duplex steel 2205 and carbon steel 1008 in two eutectic mixtures of NaCl+Na2CO3 at 700 °C and Li2CO3 + K2CO3 + Na2CO3 at 450 °C in air were studied as candidates for containment materials. A combination of optical microscopy, electrochemical measurements, SEM, EDS and EBSD techniques were employed to characterize the degradation mechanisms. Results are summarized as below.

Electrochemical measurements showed the most anodic potential values for alloy 1008 in both molten salt environments. Alloys 316 and 2205 showed more noble potential values with the same susceptibility to each molten salt. Impedance spectroscopy suggested the formation of films on the surface of the studied alloys in both molten salts which was not confirmed for alloy 2205 in chloride carbonate. And alloy 316 showed the lowest corrosion current densities in each molten salt.

Oxidation was found as the primary attack to the alloys in both molten salt environments. However, as the availability of O2 controls the degree of oxidation, both molten salts proved to slow the oxidation down when the materials were submerged.

The corroded metal morphology was grain boundary oxidation for all studied alloys in chloride carbonate at 700 °C. Corrosion in ternary carbonate at 450 °C was of uniform morphology on the surface of the alloys with no localized degradation.

Degradation of alloy 1008 in both environments involved the formation of a thick layer of iron oxide on the surface. The grain boundary oxidative attack was so severe on this alloy at 700 °C that it looked like uniform corrosion after long exposure times.

Depletion of alloying elements from grain boundaries at 700 °C contributed to intergranular attack in alloys 316 and 2205. Formation of a semi-protective oxide on the surface of alloy 316 provided more protection against chloride carbonate at 700 °C. A continuous and adherent oxide layer was also observed on the surface of this alloy at 450 °C. In alloy 2205, selective oxidation of ferrite was observed at 700 °C with no adherent film detected on the surface. The austenite phase in this alloy appeared to help protect the underlying ferrite layers at 700 °C.

Acknowledgement

This work was funded by the Australian Solar Thermal Research Initiative (ASTRI), which is supported by the Australian Government via the Australian Renewable Energy Agency (ARENA). The authors would also like to thank AINSE Ltd for providing financial assistance (Award-PGRA-2016) to enable

 

 

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work on the reported topic. The data reported in the paper were obtained at the Central Analytical Research Facility (CARF) operated by the Institute for Future Environments at Queensland University of Technology (QUT). Access to CARF was supported by generous funding from the Science and Engineering Faculty, QUT.

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Table of figures

Fig. 1 Schematic of the corrosion vessel and a sample inside the furnace. Fig. 2 (a) Open circuit potential, and (b) potentiodynamic polarization curves of samples in molten chloride carbonate at 700 °C and ternary carbonate at 450 °C. Fig. 3 Column charts of corrosion current density values in chloride carbonate at 700 °C (a) and ternary carbonate at 450 °C (b). Fig. 4 (a) Nyquist, and (b) Bode-Phase plots resulted from impedance measurements of the alloys in chloride carbonate at 700 °C, (c) Nyquist, and (d) Bode-Phase plots resulted from impedance measurements of the alloys in ternary carbonate at 450 °C. Larger Nyquist plots of CS1008, calculated equivalent circuit and fit data points are also included in (a) and (c). Fig. 5 Optical (a to c) and SEM (d to f) images of as-polished samples after 3 min exposure to the molten chloride carbonate at 700 °C. Inserted EBSD phase color map of as-received material and parallel features in (c) point to the alternate layers of ferrite/austenite in the texture of the duplex steel 2205 corroded with different rates. 

 

 

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Fig. 6 (a) Macro-graph of a SS316 sample hung over the molten chloride carbonate at 700 °C for 120 h, (b) optical image of the alloy cross-section from sidewall above the molten salt, (c) optical image of the alloy cross-section from sidewall under the salt level. Fig. 7 (a) SEM image, and (b) to (f) its corresponding EDS map analysis of an area on the top corner of SS316 above the chloride carbonate salt level and exposed to air for 120 h at 700 °C. Fig. 8 (a) SEM image, and (b) to (f) its corresponding EDS map analysis of an area on the bottom corner of SS316 submerged into the chloride carbonate salt for 120 h at 700 °C. Fig. 9 (a) Optical image of the top corner of SS316 above the ternary carbonate salt and exposed to air at 450 °C for 120 h, (b) SEM image of the red rectangular area in (a), (c) its corresponding EDS map analysis, (d) optical image of the bottom corner of SS316 under the ternary salt at 450 °C for 120 h, (e) SEM image of the red rectangular area in (d), and (f) its corresponding EDS map analysis. Fig. 10 (a) SEM image of an area on the top half of CS1008 above the chloride carbonate salt exposed to air for 120 h at 700 °C, (b) SEM-EDS map analysis corresponding to the red rectangular area in (a), (c) SEM image of an area on the bottom half of CS1008 under the chloride carbonate salt for 120 h at 700 °C, (d) SEM-EDS map analysis corresponding to the red rectangular area in (c). Fig. 11 (a) Optical image of an area on the top corner of CS1008 above the ternary carbonate salt at 450 °C for 120 h, (b) SEM-EDS map analysis corresponding to the red rectangular area in (a), (c) optical image of an area on the bottom corner of CS1008 submerged into the ternary carbonate salt at 450 °C for 120 h, (d) SEM-EDS map analysis corresponding to the red rectangular area in (c). Fig. 12 (a) Optical image from top corner of DS2205 above the chloride carbonate salt and exposed to air for 120 h at 700 °C, (b) EBSD band contrast image of the red rectangular area in (a), (c) EBSD phase and elemental analysis maps corresponding to (b), (d) optical image from bottom corner of DS2205 under the chloride carbonate salt for 120 h at 700 °C, (e) EBSD band contrast image of the red rectangular area in (d), (f) EBSD phase and elemental analysis maps corresponding to (e). Fig. 13 (a) and (b) SEM images from the side edge and top edge, respectively, on the top part of DS2205 above the chloride carbonate salt exposed to air for 120 h at 700 °C, (c) and (d) SEM images from the side edge and bottom edge, respectively, on the bottom part of DS2205 under the chloride carbonate salt for 120 h at 700 °C. Fig. 14 (a) SEM image from the top part of DS2205 above the ternary carbonate salt exposed to air at 450 °C for 120 h, (b) EDS map analysis corresponding to (a). Fig. 15 SEM image from the bottom part of DS2205 below the ternary carbonate salt at 450 °C for 120 h.