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Article

Effect of Molten Salts Composition on the Corrosion Behavior of Additively Manufactured 316L Stainless Steel for Concentrating Solar Power

by
Najib Abu-warda
*,
Sonia García-Rodríguez
,
Belén Torres
,
María Victoria Utrilla
and
Joaquín Rams
Departamento de Matemática Aplicada, Ciencia e Ingeniería de Materiales y Tecnología Electrónica, Universidad Rey Juan Carlos, Tulipán s/n, 28933 Móstoles, Spain
*
Author to whom correspondence should be addressed.
Metals 2024, 14(6), 639; https://doi.org/10.3390/met14060639
Submission received: 17 April 2024 / Revised: 22 May 2024 / Accepted: 24 May 2024 / Published: 28 May 2024
(This article belongs to the Special Issue Novel Insights and Advances in Steels and Cast Irons)
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Abstract

:
The effects of different molten salts on the corrosion resistance of laser powder bed fusion (L-PBF) 316L stainless steel was evaluated at 650 and 700 °C. The samples were characterized via XRD and SEM/EDX after high-temperature corrosion tests to evaluate the corrosion damage to the L-PBF 316L stainless steel caused by the molten salts. The presence of the salts accelerated the corrosion process, the chloride-based salts being the most aggressive ones, followed by the carbonate-based and the nitrate/nitrite-based salts, respectively. The L-PBF 316L did not react strongly with the nitrate/nitrite-based salts, but some corrosion products not found in the samples tested in the absence of salts, such as NaFeO2, were formed. LiFeO2 and LiCrO2 were identified as the main corrosion products in the samples exposed to the carbonate-based molten salts, due to the high activity of Li ions. Their growth produced the depletion of Fe and Cr elements and the formation of vacancies that acted as diffusion paths on the surface of the steel. In the samples exposed to chloride-based molten salts, the attacked area was much deeper, and the corrosion process followed an active oxidation mechanism in which a chlorine cycle is assumed to have been involved.

1. Introduction

Concentrated solar power (CSP) plants have increased in number in recent decades for decarbonizing the electricity grid due to their ability to integrate cost-effective thermal energy storage (TES) [1,2,3]. The use of TES materials that accumulate heat during sunshine periods and release it during periods of no solar irradiation has been extensive to overcome the problem of intermittency [4]. In fact, it is expected that 77% of the new CSP plants will integrate TES systems [5]. The next generation of CSP plants is expected to be more efficient in terms of energy production and this is directly related to the maximum working temperature of CSP plants and their capacity to store this energy for a time [6,7].
There are different methods employed for TES, such as latent heat in phase change materials, thermochemical, or solid sensible heat storage [8]. Sensible heat storage based on molten salts is the most used. The salt energy content increases with temperature and, when the material cools, the stored energy is released without any phase change during charge or discharge [9]. In this context, the current CSP plants use a non-eutectic salt mixture of 60 wt.% NaNO3 and 40 wt.% KNO3 [10]. This commercial salt, commonly known as solar salt, has high heat capacity and low price; however, it also has low thermal stability limits. Its use supposes that the maximum operating temperature is near 550 °C [11]. In this context, other salt mixtures have been considered to overcome this maximum operation limit. Particularly, carbonate-based 32K2CO3-35Li2CO3-33Na2CO3 salt (in wt.%) has a higher thermal stability temperature, but the price of Li2CO3 limits its use [12,13]. Other TES materials based on chloride salts, such as 23.9KCl-7.5NaCl-68.6ZnCl2 or 17.8KCl-14NaCl-68.2ZnCl2 have been proposed because they meet the requirements of high thermal stability limits and low price [14]. In addition, chloride salts, such as NaCl and KCl, are available in large amounts, and when mixed with ZnCl2 or MgCl2, eutectic mixtures with low melting points and high stability limits are achieved. However, despite these advantages, chloride-based TES materials have not yet been implemented due to metallurgical problems related to their high corrosiveness [15,16]. For this reason, more efforts are needed to understand the effect of these new TES materials on the high-temperature corrosion behavior of metal alloys used as structural components in CSP plants.
In this context, austenitic stainless steels, and specifically, 316L stainless steel, are the most used structural materials in many industries, mainly due to their good corrosion resistance, which is attributed to the addition of Cr and Mo [17,18,19], and their good mechanical properties at high temperatures [20]. Moreover, its good oxidation resistance makes 316L stainless steel an excellent candidate to be used in CSP plants and high-temperature parts such as pipelines or hot tanks manufactured with stainless steel [21,22,23].
In addition, the manufacturing process also plays an important role in the corrosion behavior of 316L ss. In a previous study, the influence of the manufacturing process on corrosion resistance was evaluated in the presence of 17.8KCl-14NaCl-68.2ZnCl2 molten salt [24]. The 316L manufactured by laser powder bed fusion (L-PBF) provided more corrosion resistance than the wrought one, due to the presence of more pathways for Cr diffusion associated with a higher number of defects in the microstructure of the L-PBF samples, facilitating the formation of the protective Cr2O3 scale. In this context, additive manufacturing (AM) techniques have attracted great interest in recent years to produce components for different industry sectors.
AM encompasses a set of manufacturing techniques based on the production of components through adding material layer by layer from a three-dimensional (3D) model [25]. In the field of metals, L-PBF, also known as selective laser melting (SLM), is one of the most used AM techniques [26,27,28]. Among other advantages, these techniques stand out due to the possibility of manufacturing components with more complex geometries than those obtained with casting processes [29].
The design flexibility of AM allows the manufacturing of pieces as single components, avoiding the need to weld different parts [30]. These new AM techniques could be implemented in valves, complex pipes, pumps, or complex heat exchangers in the energy sector [31]. However, the corrosion behavior of alloys manufactured through L-PBF in aggressive environments has not been studied. In this context, this work presents a comparative analysis of the corrosiveness of different TES molten salts, analyzing their effect on L-PBF 316L stainless steel. The results of this research provide helpful information about both the use of these molten salts as TES materials in the next generation of CSP plants and their effect on 316L stainless steel manufactured through L-PBF.

2. Materials and Methods

2.1. Base Material

The austenitic 316L ss produced via L-PBF with a size of 15 × 15 × 3 mm3 was manufactured in a vertical orientation and was provided by Idonial (Madrid, Spain). The additive manufacturing system was a fiber laser M280 400W (EOS, Krailling, Germany). The samples were manufactured with 1070 nm laser beam wavelength, 0.07 mm laser spot diameter, and 195 W power. In addition, 1083 mm/s scanning speed, 20 µm layer height, and 90 µm hatching distance were used. The composition of the powders used in the L-PBF process was 17.0 Cr, 13.1 Ni, 2.7 Mo, 1.3 Mn, 1.2 Si, ≤0.1 O, ≤0.03 C, and Fe (rest) (in wt.%), and the particle size of the powder was in the range of 20–50 μm.

2.2. Corrosion Tests

Hot corrosion tests were performed to evaluate the effect of salt composition on the 316L stainless steel without any surface treatment. In this context, the as-built samples had an average surface roughness (Ra) of 2.8 ± 0.6 µm. Different salt mixtures were used for this purpose, and their composition and nomenclature are shown in Table 1. The first batch of samples was tested without any salt mixture and was used for comparative purposes. The second batch of samples was exposed to a nitrate/nitrite-based salt mixture, which was selected because it is the most implemented molten salt of the four analyzed and was employed as a reference salt in this work [32]. The KNaNOx combinations are unstable at the temperature used in this study, but they have been included as a common reference and to study conditions in which there is a punctual overheating of the system. The third batch was exposed to a carbonate-based salt mixture [33], and the last two batches of samples were exposed to two different chloride-based salt mixtures, one rich in ZnCl2 [34] and the other one rich in MgCl2 [35]. The latter three molten salts were investigated because they are promising TES materials for use in CSP plants.
High-temperature corrosion tests were performed in a dry air atmosphere to simulate a punctual failure in the inertization system of a plant using LT 5/12/P330 furnace (Nabertherm GmbH, Lilienthal, Germany), following the standard ISO/FDIS 17224:2014 for high-temperature corrosion tests [36]. Figure 1 shows a scheme of the testing method employed for the corrosion tests; two samples were used for each studied condition. A short exposure time (48 h) was selected, also to simulate the punctual failure in the inertization system. In addition, two temperatures (650 and 700 °C) were employed. These temperatures were selected because the working temperature is expected to rise in the next generation of CSP plants and they are near the thermal stability limits of the employed carbonate-based and chloride-based salts, assuming that this temperature range is not ideal for the nitrate/nitrite-based salts as it is higher than their stability limit.
The corrosion rate produced on the L-PBF 316L ss exposed to the different molten salts was quantified using metal loss distribution, via a dimensional metrology analysis. The analyses were performed by measuring cross-section thickness of the uncorroded steel (fifty measurements in each sample) at equidistant distances before (e0) and after (ef) the corrosion test, as Figure 1 represents, using DMR LAS V4 light microscope (Leica, Wetzlar, Germany) with Image-Pro Plus software (version 7.0, Leica, Wetzlar, Germany). In this manner, the entire cross-section of the sample was assessed. The difference between e0 and ef was considered as metal loss. For the measurement of the thickness after corrosion tests (ef), only the thickness of the uncorroded steel was considered. For this reason, it was not necessary to clean the residual salts and corrosion products from the sample’s surface after the high-temperature corrosion tests. The data sets were re-ordered (from greatest to least metal loss) and plotted as metal loss (in µm) vs. cumulative probability (%). This type of plot indicates the probability (in %) of a certain degree of damage being observed.

2.3. Samples Characterization

The as-built 316L austenitic stainless steel was first characterized by optical microscopy (OM) to analyze its microstructure. The steel was ground and polished and then electrochemically etched with oxalic acid to reveal the microstructure. A LEICA/DMR optical microscope equipped with Image-Pro Plus analysis software (version 7.0, Leica, Wetzlar, Germany) was employed for this purpose. The cross-sections of the sample were analyzed using a scanning electron microscope (SEM, Hitachi S-3400 N, Hitachi High-Technologies, Tokyo, Japan) equipped with a XFlash 5010 energy dispersive X-ray spectrometer (EDS, Bruker, Germany).
The X-ray diffraction (XDR) technique was used to identify the corrosion products. A Bruker AXS D8 Advance diffractometer was employed using CuKα1 radiation (λ = 1.54 nm) in θ–2θ mode. Diffraction patterns were recorded at an angular interval of 20 to 90° and indexed using the PDF-4 database of ICDD.

3. Results and Discussion

3.1. Characterization of the Starting Materials

The composition of the 316L ss was 17.0 Cr, 13.1 Ni, 2.7 Mo, 1.3 Mn, 1.2 Si, and bal. Fe (in wt.%). XRD spectrum of the 316L stainless steel manufactured through L-PBF is shown in Figure 2a. The pattern shows that the sample was characterized only by the FCC austenite phase (00-003-1209, ICDD).
Figure 2b shows an optical micrograph of the main plane (ZY) of the steel. Its macroscale structure was determined by the L-PBF process, which depended on the laser-scanning pattern. In this direction, a layered microstructure can be observed, composed of a network of melt pools formed from the overlap of the laser-scanning patterns. It can also be observed that large columnar austenitic grains formed within the melt pools, typical of this manufacturing process. These were formed in the direction of the thermal gradient (Z direction) through the melt pool boundaries [37]. Figure 2c illustrates the microstructure of the alloy at higher magnification. The cellular microstructure that characterizes the melt pools is shown. This microstructure is related to the manufacturing process, where a high cooling speed induced non-equilibrium solidification conditions [38]. These cells had different morphologies, equiaxed–polygonal or elongated, depending on their growth direction.

3.2. Measurement of Metal Damage

The use of aggressive salts usually produces the growth of corrosion products, which can be lost or gained by evaporation or sublimation of gaseous species. These species can produce slight distortions in mass change measurements and, therefore, the corrosion and damage suffered by the 316L stainless steel was analyzed using dimensional metrology analysis. The metal damage produced in the 316L L-PBF stainless steel after its exposure to the different salt mixtures in terms of metal loss (in µm) vs. cumulative probability (in %) is represented in Figure 3a.
As expected, in all the tested salt mixtures, the steel suffered more damage at 700 than at 650 °C. However, the effect of temperature was negligible in the samples tested without any salt mixture, and for both temperatures, the metal loss was very low (~25 µm). Nevertheless, different trends can be observed in Figure 3a. The plots of the samples exposed to high temperatures without any molten salts (oxidation test) were straight at both temperatures and metal loss values were almost constant, indicating that no localized attack was produced, and that the oxidation process was continuous in all the samples, favoring the grown of an oxide layer with a similar thickness along the entire surface. The same trend was observed in the samples exposed to KNaNOx molten salts, probably because at these temperatures, KNaNOx is unstable and breaks down easily. However, the samples exposed to LiNaKCO3 molten salts at 650 and 700 °C suffered a more heterogeneous attack. As a representative example, results for the samples tested at 700 °C showed that approximately 20% of the surface underwent more attack compared with the other 80% of the surface, which suffered less attack. The same trend was observed in the samples exposed to NaKZnCl and NaKMgCl at 700 °C, where almost 10% and 15% of sample surface, respectively, underwent more corrosion. The localized attack shown in these samples can be attributed to chloride’s aggressiveness, which accelerated the corrosion process in some areas.
Figure 3b shows the average metal loss data reported from Figure 3a, for the L-PBF 316L stainless steel after its exposure to the different molten salts to classify these based on their aggressiveness. As expected, the most aggressive salts were those rich in chlorides, particularly NaKZnCl, followed by NaKMgCl. At 700 °C, these salts increased the metal loss twenty and fifteen times compared with the samples exposed to high temperatures without any molten salts. KNaNOx and LiNaKCO3 molten salts produced lower damage in the stainless steel and increased the metal loss four and five times at 700 °C compared with the samples exposed to high temperatures without any molten salts. In addition, the metal loss increased when the exposure temperature was raised to 700 °C in all the molten salts tested.

3.3. Corrosion Products Characterization

3.3.1. Visual Analysis

The macrographs presented in Figure 4 show the surfaces of the samples after their exposure to the different TES materials. As can be seen, the composition of the salt had more influence than the exposure temperature because according to their degree of damage, there were no relevant differences between the samples tested at 650 and 700 °C. However, the exposure to the molten salts produced huge increases in oxides and corrosion product content on their surfaces. The steels exposed to high temperatures without any salts (oxidation tests) still showed their metal surface, while the samples exposed to nitrate/nitrite-based salts and carbonate-based salts no longer showed their metal surface (Figure 4). In these samples, a continuous layer of oxides/corrosion products was formed, but these still adhered to the sample. Otherwise, the 316L stainless steel samples exposed to chloride-based salts were highly attacked and corroded. The oxides and corrosion products exerted extreme delamination in these samples. These results are in accordance with those reported in Figure 3, which also confirmed not only the stronger influence of molten salt presence compared with difference in temperature in the corrosion process but also the order of aggressiveness of the studied TES materials. The nitrate/nitrite-based and carbonate-based molten salts were much less aggressive than the chloride-based molten salts.

3.3.2. XRD Characterization

Figure 5 shows the XRD patterns of the tested samples at 650 °C, Figure 5a, and 700 °C, Figure 5b, and Table 2 shows a list of the XRD-identified species to facilitate the analysis of the results. According to the visual examination shown in Figure 4 of the samples exposed to high temperatures in the absence of molten salts, the prominent peaks were those of austenite γ-Fe; however, some small peaks of hematite (Fe2O3) and spinel ((Mn,Fe)Cr2O4) oxides were also identified at both tested temperatures.
In the samples exposed to the KNaNOx salt mixture, the main peaks were associated with NaFeO2 [39]; however, some oxides that adhered well to the surface (Figure 4), such as Fe2O3, Fe3O4, (Mn,Fe)Cr2O4, and NiCr2O4, were also identified. Peaks of γ-Fe were identified at 650 °C but not at 700 °C. This is because of the presence of a thicker layer of oxides and corrosion products in the sample tested at 700 °C (see Figure 3), which would prevent the identification of the γ-Fe phase via XRD.
The XRD patterns of the 316L stainless steel tested under LiNaKCO3 molten salts at 650 and 700 °C revealed that at both temperatures, the corrosion products formed on the samples’ surfaces were LiFeO2, LiCrO2, and Fe3O4. These were the main corrosion products, and Luo et al. [40] also identified these peaks when they studied the corrosion of the same stainless steel by LiNaK carbonate molten salt at 600 and 700 °C. González-Fernández et al. [41] analyzed the corrosion products formed in 310 ss (24.9% Cr) at 600 °C for a longer exposure time (600 h) in dynamic conditions and also found the presence of Fe3O4, LiCrO2, and Li(Fe,Ni)O2 as the main corrosion products. This means that the short exposure time was enough to produce the lithiation reaction of the grown oxides. González-Fernández et al. also found the presence of MnNi6O8 and K2NiO2; in contrast, however, in our case, these compounds were not identified at lower exposure times.
In the samples exposed to the NaKMgCl and NaKZnCl chloride-based TES materials, peaks of γ-Fe were also identified; however, in this case, the peaks appeared due to the high delamination suffered by the oxide coating, as Figure 4 previously revealed, that exposed the uncorroded material to the surface. In the samples exposed to NaKZnCl molten salt, peaks of Cr2O3, ZnCr2O4, and NiCr2O4 were also identified at both temperatures. However, in the NaKMgCl molten salt, the most relevant result was the presence of MgO, which other authors also identified as a corrosion product in Ni-rich alloys [42,43]. In this context, Gong et al. [44] analyzed the compatibility of Fe-based alloys with purified molten MgCl2-KCl-NaCl salt at the same temperature (700 °C) for longer periods and they also found the formation of MgO as the main corrosion product after 1000 h.

3.3.3. Cross-Section Characterization

Figure 6 shows the SEM analysis of the cross-section of the L-PBF 316L ss exposed to 650 and 700 °C without any molten salts. The microstructure of the grown oxide layer was extensively characterized in a previous study [24]. It was formed of a homogeneous coating of ~6 and ~9 µm thickness on the samples tested at 650 and 700 °C, respectively, and it was composed from the inner to the top surface of MnCr2O4, Fe2O3, and FeCr2O4. In addition, the oxide layer adhered well to the substrate; no delamination or cracks were observed in the substrate–oxide interface, as Figure 6c shows at higher magnification for the L-PBF 316L stainless steel exposed at 700 °C for 48 h.
Figure 7 shows the SEM micrographs of the L-PBF 316L stainless steel exposed to the KNaNOx salt at 650 °C (Figure 7a,b) and 700 °C (Figure 7c–e). The main difference observed between the two test temperatures was the thickness of the grown oxides coating, which was of ~8 µm and ~20 µm thickness at 650 and 700 °C, respectively. Figure 7a shows that there was not a great reaction of the steel with the salt mixture, probably because at this range of temperatures, KNaNOx salt is unstable and breaks down easily. Nevertheless, some important differences were seen compared with the samples tested in that absence of salt and different layers were observed in the grown coating, as the EDX map of Figure 7b shows. The inner layers were rich in Mn, Cr, and Ni and, according to the XRD results, (Mn,Fe)Cr2O4 and NiCr2O4 oxides grew. The intermediate and outer layers had more content of Fe and Na, and they may have been composed, respectively, of Fe3O4, Fe2O3, and NaFeO2. The formation of iron oxides was mainly initiated by the reaction of iron with the oxygen of the atmosphere and NaFeO2 was produced due to the reaction of Na2O with these iron oxides. Ahmed [45] also found NaFeO2 on 304 stainless steel after exposure to eutectic KNO3-NaNO3 and he explained the presence of a sodium iron oxide phase on the top region of the bulk corrosion scale with the following reactions (1) and (2):
2NaNO2 ↔ Na2O + NO + NO2
Na2O + Fe2O3 ↔ 2NaFeO2
In addition, the presence of K2CrO4, which was not identified with XRD, would explain the presence of K in the EDX map, as it has been observed in Fe-Cr alloys in presence of KCl [46]. Other authors, such as Sandoval-Amador et al. [17] and Zhang et al. [47], analyzed the corrosion of conventionally manufactured 316L ss in presence of molten NaNO3-KNO3-NaNO2 and NaNO3-KNO3 at lower temperatures (450 to 550 °C for 150 h), and they identified iron and spinel oxides as the main corrosion products. Gomes et al. [48] analyzed the behavior of the same stainless steel in a mixture of 60 wt.% NaNO3 and 40 wt.% KNO3 at 550 °C, and again, they found Fe2O3, Fe3O4, and FeCr2O4 as the main scale products. In their study, the thickness of the scale layer formed on AISI 316L was 6.9 ± 2.1 µm, like that reported in the present study; however, while the exposure time was longer (3000 h), the exposure temperature was 100 °C lower.
Figure 7d shows the surface of the L-PBF 316L ss after its exposure to the KNaNOx at 700 °C. The red arrows in Figure 7d indicate the accumulation of oxides in some areas and, at higher magnification, the morphology of the oxides forming small octahedral crystals of ~5 µm can be differentiated (Figure 7e). The EDX analysis performed in this area (in Figure 7d) revealed the presence of only Fe, Na, and O, which are associated with NaFeO2 and Fe-rich oxides.
Figure 8 shows the SEM micrographs of the L-PBF 316L stainless steel exposed to carbonate-based (LiNaKCO3) molten salts at 650 °C (Figure 8a–c) and 700 °C (Figure 8d–g) for 48 h. The continuous oxide layer shown previously using the KNaNOx salt was not observed, but the degree of corrosion was slightly superior at 650 °C, Figure 8a. In addition, Figure 8b shows at higher magnification that the area near the surface was porous and full of cavities and holes (marked with red arrows in Figure 8a,b). These vacancies were left by the depletion of some elements, thereby forming diffusion paths [49]. The formation of these diffusion paths during the high-temperature corrosion process can be attributed to unbalanced diffusional processes known as the Kirkendall effect. In this area, the main elements of the alloy (Fe and Cr) were not present, as the EDX mappings of Figure 8c revealed. In contrast, this area was enriched in Ni and Mo. Otherwise, the presence of Li could not be detected through EDX (it is important to consider the low sensitivity of EDX to light elements), but some Li-rich phases, such as LiFeO2 and LiCrO2, were detected with XRD, as shown in Figure 5. The high activity of Li ions in the molten salt produced their reaction with Fe and Cr-rich oxides, forming Li-rich compounds (LiFeO2 and LiCrO2) following reactions (3) and (4). LiCrO2 has a higher solubility than LiFeO2 in the molten carbonate salt; therefore, a stable LiFeO2 layer gradually formed on the top.
Fe2O3 + O2− + 2Li+ ↔ 2LiFeO2
Cr2O3 + O2− + 2Li+ ↔ 2LiCrO2
Keijzer et al. [50] analyzed the presence of Li-rich compounds on AISI 304 stainless steel, showing an outermost Fe-rich layer followed by a Cr-rich layer next to the steel. The LiCrO2 layer acted as a strong diffusion barrier for metal ions on top of the oxide layer, while the LiFeO2 layer above helped to protect the Cr-rich layer from chromate formation. However, at the same time, the formation of these layers produced high stress in both the metal and the oxide layers and led to cracking. This probably produced the delamination of the metal at 700 °C shown in Figure 8d,e under high magnification.
Sarvghad et al. [51] tested the same molten salts with conventional 316 stainless steel and reported that the steel did not suffer severe damage and no localized dealloying was produced. However, they tested the steel at lower temperatures (450 °C). In addition, Luo et al. [40] employed the same TES material in 316L ss. They reported that the corrosion process was mainly divided into four steps: (i) oxidation of the alloy, (ii) lithiation reaction of the oxides, (iii) formation of a dual-structured corrosion scale, and (iv) spallation of corrosion products. These steps may also be valid for the present work, but probably with changes in the kinetic reactions attributed to differences in the alloy’s fine sub-granular cellular microstructure caused by the manufacturing process. Figure 8f represents the surface of the 316L exposed to LiNaKCO3 at 700 °C, and the Li-rich corrosion products and the oxides forming octahedral crystals are shown at higher magnification in Figure 8g.
Figure 9a shows the SEM analysis of the L-PBF 316L ss after its corrosion by NaKMgCl molten salts at 650 °C. In this case, the steel suffered severe damage, and the affected area was much deeper than with the previously analyzed molten salts. The steel was full of corrosion pathways, and different layers of corrosion products were detected, the outer one being formed by oxides. The corrosion process followed an active oxidation mechanism in which a chlorine cycle is assumed to have been involved due to the molten salt’s presence, where MgO originated from the spontaneous decomposition of MgOHCl, which was generated from the hydrolysis of hydrated MgCl2 as MgCl2·H2O, following reaction (5). Then, chlorine reacted with metal elements to produce volatile metal chlorides via reaction (6), escaping the corrosion system [52]. In this context, some alloying elements like Cr, which has lower standard electrode potential in molten chlorides compared with Fe or Ni, could not avoid the corrosion process [53]. In this system, volatiles like CrCl4 and oxide precipitates like MgO are considered the main corrosion products based on the thermodynamic calculation [54].
MgCl2 + H2O ↔ MgO + 2HCl
nCl2 + 2M ↔ 2MCln
Figure 9b shows the exposed L-PBF 316L stainless steel surface at 700 °C. This micrograph shows the inner and outer layers of corrosion products because the outer layer spalled off in some regions. The border of these two different scales can be appreciated; punctual EDX analysis was performed on each region. The inner layer, marked in Figure 9b as (1), was rich in Fe (31%), Ni (36%), and O (20%); however, the presence of other elements such as Cr (2%) or Mg (no signal) was residual. In contrast, the outer layer of corrosion products, identified in Figure 9b as (2), was higher in Cr (45%), O (32%), and Mg (6%), but it had a lower amount of Fe (15%) and no evidence of Ni. According to these results, MgO, as also identified by Wang et al. [43], MgCr2O4, and other Cr-rich oxides were formed above the layer of corrosion products.
Figure 10 represents SEM micrographs of the L-PBF 316L ss after its exposure to NaKZnCl molten salts at 700 °C. The behavior of the steel under this TES material was also analyzed in a previous study [24]. Figure 10 shows how the corrosion progressed through the fine sub-grains characteristic of the L-PBF manufacturing process, without generating preferential paths for diffusion of chlorides. In this context, a high number of pathways associated with the high number of grain boundaries in the microstructure of L-PBF parts were available for the diffusion of Cr to the surface and this facilitated the enrichment in Cr near the surface and the formation of a Cr2O3 scale. However, the presence of the NaKZnCl molten salts induced an active oxidation mechanism, which promoted the chlorides’ formation on the steel surface through the reaction of the grown Cr2O3 scale with the ZnCl2 salt (see reaction (7)) [55] and accelerated the corrosion process through the formation of MClx products, as reactions (8) and (9) show. The chlorides diffused through the Fe-rich oxides to the oxide–metal interface, leading to the formation of solid iron chloride. However, the high volatility of the iron chloride provoked its diffusion backwards through the non-protective Fe-rich oxide layer.
ZnCl2 + Cr2O3 + 1/2O2 ↔ ZnCr2O4 + Cl2
Fe + Cl2 ↔ FeCl2
Cr + Cl2 ↔ CrCl2
Once a sufficient oxygen partial pressure was reached, the chlorides destabilized and formed Cr2O3 or Fe2O3, while Cl2 regenerated (reactions (10) and (11)). These reaction steps were continually repeated, causing severe damage in the alloy. In that case, the large amount of oxides and corrosion products on sample surface could have prevented the detection via XRD of the Fe3O4 phase identified in the samples exposed to NaKMgCl molten salts, but this does not mean that it was not present where Fe2O3 had grown on top of the Fe3O4 layer.
2FeCl2 + 3/2O2 ↔ Fe2O3 + 2Cl2
2CrCl2 + 3/2O2 ↔ Cr2O3 + 2Cl2
The high corrosion rates observed for the L-PBF 316L stainless steel, particularly in those samples exposed to the Cl-rich salt mixtures, highlight the importance of other alternatives, as the use of protective coatings [56,57], surface modification methods [58,59,60], the purification of the salts to avoid the presence of contaminants that can accelerate the corrosion process [44,61,62,63], or the use of corrosion inhibitors [64].
In addition, further investigations should be performed using an inert gas in absence of an air atmosphere, because hydrolysis products, e.g., HCl and MgOHCl, can accelerate the corrosion process, particularly in the presence of chloride-based salts. From this research, we can conclude that this alloy cannot be a feasible option to consider in the design of a plant which will use chloride-based molten salts at temperatures above 650 °C or, at least, not under the studied conditions where a failure in the inertization system was simulated.

4. Conclusions

The behavior of the L-PBF 316L stainless steel was more affected by the compositions of the investigated molten salts than by an increase in temperature from 650 to 700 °C, the most aggressive salts being those rich in chlorides.
A homogeneous and well-adhered oxide layer was formed on the surface of L-PBF 316L stainless steel exposed to high temperatures without any molten salts. In addition, there was the L-PBF 316L stainless steel did not react strongly with the KNaNOx salt, partially because of the degradation of the salt at the temperatures used. In these samples, the inner layers of the oxide coating were formed of (Mn,Fe)Cr2O4 and NiCr2O4 and the outer layer was composed of NaFeO2, Fe3O4, and Fe2O3.
The high activity of Li ions in the LiNaKCO3 molten salts produced their reaction with Fe and Cr-rich oxides, forming LiFeO2 and LiCrO2 as the main corrosion products, and their growth produced the depletion of Fe and Cr elements and the formation of vacancies which acted as diffusion paths in the surface of the steel.
In the samples exposed to chloride-based molten salts, the corroded area was much deeper and the corrosion process followed an active oxidation mechanism where chlorine reacted with metal elements to produce volatile metal chlorides which accelerated the corrosion process.

Author Contributions

Conceptualization, N.A.-w., S.G.-R. and M.V.U.; methodology, N.A.-w. and S.G.-R.; formal analysis, N.A.-w. and S.G.-R.; investigation, N.A.-w. and S.G.-R.; writing—original draft preparation, N.A.-w. and S.G.-R.; writing—review and editing, B.T., M.V.U. and J.R.; supervision, M.V.U. and J.R.; project administration, B.T. and J.R.; funding acquisition, B.T. and J.R. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by Ministerio de Ciencia e Innovación (PID2021-123891OB-I00, PID2021-124341OB-C21).

Data Availability Statement

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also form part of an ongoing study.

Conflicts of Interest

The authors declare no conflicts of interest.

References

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Figure 1. Scheme of the testing method employed for corrosion tests.
Figure 1. Scheme of the testing method employed for corrosion tests.
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Figure 2. (a) X-ray diffraction analysis, (b) optical microscopy characterization, and (c) SEM micrograph of the as-built L-PBF 316L stainless steel sample.
Figure 2. (a) X-ray diffraction analysis, (b) optical microscopy characterization, and (c) SEM micrograph of the as-built L-PBF 316L stainless steel sample.
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Figure 3. High-temperature corrosion behavior of L-PBF 316L stainless steel in the presence of different TES molten salts at 650 and 700 °C for 48 h: (a) metal loss (in µm) vs. cumulative probability (in %) and (b) average metal loss results (data of NaKZnCl samples from Ref [24]).
Figure 3. High-temperature corrosion behavior of L-PBF 316L stainless steel in the presence of different TES molten salts at 650 and 700 °C for 48 h: (a) metal loss (in µm) vs. cumulative probability (in %) and (b) average metal loss results (data of NaKZnCl samples from Ref [24]).
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Figure 4. Surface morphology of L-PBF 316L stainless steel after its exposure to high-temperature corrosion tests at 650 and 700 °C for 48 h in the presence of different TES materials.
Figure 4. Surface morphology of L-PBF 316L stainless steel after its exposure to high-temperature corrosion tests at 650 and 700 °C for 48 h in the presence of different TES materials.
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Figure 5. XRD patterns of the L-PBF 316L stainless steel after its exposure to different TES materials at (a) 650 and (b) 700 °C for 48 h.
Figure 5. XRD patterns of the L-PBF 316L stainless steel after its exposure to different TES materials at (a) 650 and (b) 700 °C for 48 h.
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Figure 6. SEM characterization of the L-PBF 316L ss exposed to dry air atmosphere at (a) 650 °C and (b,c) 700 °C for 48 h at different magnification (Adapted from Ref. [24]).
Figure 6. SEM characterization of the L-PBF 316L ss exposed to dry air atmosphere at (a) 650 °C and (b,c) 700 °C for 48 h at different magnification (Adapted from Ref. [24]).
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Figure 7. SEM and EDX characterization of the L-PBF 316L ss exposed to the KNaNOx: (a,b) cross-sectional analysis of samples exposed at 650 °C and (c–e) cross-sectional and superficial analysis of samples exposed at 700 °C for 48 h.
Figure 7. SEM and EDX characterization of the L-PBF 316L ss exposed to the KNaNOx: (a,b) cross-sectional analysis of samples exposed at 650 °C and (c–e) cross-sectional and superficial analysis of samples exposed at 700 °C for 48 h.
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Figure 8. Cross-sectional and superficial SEM/EDX characterization of the L-PBF 316L ss after its exposure to the LiNaKCO3 TES material at (a–c) 650 °C and (d–g) 700 °C for 48 h.
Figure 8. Cross-sectional and superficial SEM/EDX characterization of the L-PBF 316L ss after its exposure to the LiNaKCO3 TES material at (a–c) 650 °C and (d–g) 700 °C for 48 h.
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Figure 9. SEM/EDX characterization of the L-PBF 316L ss exposed to the NaKMgCl TES material for 48 h at (a) 650 °C and (b) 700 °C.
Figure 9. SEM/EDX characterization of the L-PBF 316L ss exposed to the NaKMgCl TES material for 48 h at (a) 650 °C and (b) 700 °C.
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Figure 10. Cross-sectional SEM analysis of the L-PBF 316L stainless steel exposed to the NaKZnCl TES material at 700 °C for 48 h at (a) ×300 and (b) ×2000 magnification.
Figure 10. Cross-sectional SEM analysis of the L-PBF 316L stainless steel exposed to the NaKZnCl TES material at 700 °C for 48 h at (a) ×300 and (b) ×2000 magnification.
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Table 1. Molten salts properties, composition (in wt.%), and nomenclature.
Table 1. Molten salts properties, composition (in wt.%), and nomenclature.
Molten SaltNo SaltsNitrates/NitriteCarbonatesZn-ChloridesMg-Chlorides
NomenclatureOxidationKNaNOxLiNaKCO3NaKZnClNaKMgCl
Melting point (°C)-142397204380
Stability limit (°C)-535>650850>800
KNO3-53.0---
NaNO3-7.0---
NaNO2-40.0---
K2CO3--32.0--
Li2CO3--35.0--
Na2CO3--33.0--
KCl---23.917.8
NaCl---7.514.0
ZnCl2---68.6-
MgCl2----68.2
Table 2. XRD-identified species on the L-PBF 316L stainless steel after high-temperature corrosion tests in the presence of different salt mixtures at 650 and 700 °C for 48 h. The dominant peaks are marked in bold for each test.
Table 2. XRD-identified species on the L-PBF 316L stainless steel after high-temperature corrosion tests in the presence of different salt mixtures at 650 and 700 °C for 48 h. The dominant peaks are marked in bold for each test.
Steel/DepositOxidationKNaNOxLiNaKCO3NaKMgClNaKZnCl
650 °Cγ-Fe
Fe2O3
(Mn,Fe)Cr2O4		NaFeO2
Fe2O3
Fe3O4
NiCr2O4
(Mn,Fe)Cr2O4
γ-Fe		Fe3O4
LiFeO2
LiCrO2		Fe3O4
MgO
γ-Fe		Cr2O3
(Mn,Fe)Cr2O4
NiCr2O4
ZnCr2O4
Fe2O3
γ-Fe
NaCl
700 °Cγ-Fe
Fe2O3
(Mn,Fe)Cr2O4		NaFeO2
Fe2O3
Fe3O4
NiCr2O4
(Mn,Fe)Cr2O4		Fe3O4
LiFeO2
LiCrO2		Fe3O4
MgO
γ-Fe		Cr2O3
(Mn,Fe)Cr2O4
NiCr2O4
ZnCr2O4
Fe2O3
γ-Fe
NaCl
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Abu-warda, N.; García-Rodríguez, S.; Torres, B.; Utrilla, M.V.; Rams, J. Effect of Molten Salts Composition on the Corrosion Behavior of Additively Manufactured 316L Stainless Steel for Concentrating Solar Power. Metals 2024, 14, 639. https://doi.org/10.3390/met14060639

AMA Style

Abu-warda N, García-Rodríguez S, Torres B, Utrilla MV, Rams J. Effect of Molten Salts Composition on the Corrosion Behavior of Additively Manufactured 316L Stainless Steel for Concentrating Solar Power. Metals. 2024; 14(6):639. https://doi.org/10.3390/met14060639

Chicago/Turabian Style

Abu-warda, Najib, Sonia García-Rodríguez, Belén Torres, María Victoria Utrilla, and Joaquín Rams. 2024. "Effect of Molten Salts Composition on the Corrosion Behavior of Additively Manufactured 316L Stainless Steel for Concentrating Solar Power" Metals 14, no. 6: 639. https://doi.org/10.3390/met14060639

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