High temperature oxidation of steels in CO2 containing impurities

Journal Pre-proof High temperature oxidation of steels in CO2 containing impurities ¨ Richard P. Oleksak, Joseph H. Tylczak, Gordon R. Holcomb, Omer ˘ N. Dogan

PII:

S0010-938X(19)31109-6

DOI:

https://doi.org/10.1016/j.corsci.2019.108316

Reference:

CS 108316

To appear in:

Corrosion Science

Received Date:

28 May 2019

Revised Date:

24 October 2019

Accepted Date:

29 October 2019

¨ ˘ Please cite this article as: Oleksak RP, Tylczak JH, Holcomb GR, Dogan ON, High temperature oxidation of steels in CO2 containing impurities, Corrosion Science (2019), doi: https://doi.org/10.1016/j.corsci.2019.108316

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High temperature oxidation of steels in CO2 containing impurities

Richard P. Oleksak1,2,*, Joseph H. Tylczak1, Gordon R. Holcomb1, Ömer N. Doğan1

National Energy Technology Laboratory, 1450 Queen Ave SW, Albany, OR 97321-2198, USA

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Leidos Research Support Team, 1450 Queen Ave SW, Albany, OR 97321-2198, USA

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*Corresponding Author Email: [email protected]

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Phone: 1-541-918-4538

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1450 Queen Ave SW Albany, OR 97321-2198

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Highlights

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USA

Several steels were exposed to 95% CO2, 4% H2O, 1% O2 with/without 0.1% SO2 at 1 atm and 550 °C

Low-Cr steels formed thick Fe-rich oxide scales regardless of SO2

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High-Cr steels formed thin Cr-rich oxide scales in the absence of SO2

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SO2 caused breakdown in Cr-rich oxide growth due to S in the scale

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The evolving microstructure of the alloy sub-surface influenced the oxidation reaction

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Abstract 1

Future power systems require steels that resist corrosion in high temperature CO2 containing combustion products/impurities. Here we evaluate the corrosion behavior of several commercial steels exposed to 95%CO2-4%H2O-1%O2, with/without 0.1% SO2, at 550 °C and 1 atm. High-Cr steels formed thin Cr-rich oxides and low-Cr steels thick Fe-rich oxides during SO2-free exposures. Additions of SO2 had little effect on low-Cr steels but enhanced corrosion of high-Cr steels, where S within the chromia scale led to its failure and subsequent Fe-rich oxide nodule growth. The

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results are rationalized by considering the thermodynamic and kinetic factors controlling the

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reaction.

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Keywords: A. Stainless steel, C. High temperature corrosion, C. Oxidation, C. Carburization, C.

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Sulfidation, Supercritical CO2 power cycle, Oxyfuel combustion

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1. Introduction

Emerging power system technologies require structural materials that

resist

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oxidation/corrosion in CO2-rich gases at high temperatures (>500 °C). For example, oxyfuel combustion (the process of burning fuel in pure O2 rather than air) is being developed as a way to facilitate CO2 capture by eliminating the need for N2 separation of the exhaust gas [1]. Metallic components in an oxy-fired system are exposed to gases that contain significantly higher quantities

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of CO2 (≈60%) compared to an air-fired system (≈15%). The balance of the gas consists of H2O, O2, and potentially SO2, depending on the fuel and separation processes employed [2, 3]. Similar compositions are expected in next-generation direct-fired supercritical CO2 (sCO2) based power cycles, which offer the potential for significantly improved conversion efficiencies relative to current systems [4]. In a direct-fired (semi-open) sCO2 power cycle, the CO2-rich stream produced by an oxy-combustion is itself utilized as the working fluid [5]. Thus, metallic components are 2

exposed to a high pressure, CO2-rich fluid containing the same combustion impurities as described above. Due to their relatively low cost, Fe-Cr based steels are preferred materials for construction of current and future power systems up to the temperatures dictated by the long-term mechanical strength of the steel—usually about 650 °C. Therefore, candidate steels must also resist oxidation during long-term exposure to these conditions. As such, the high temperature oxidation of steels

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in CO2 has been studied for many years. For example, research dating to the 1960s focused on

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corrosion issues encountered with steels used in gas cooled nuclear reactors, where CO2 is the primary coolant [6-11]. More recently, the topic of steel oxidation in pure CO2 has been revisited

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by many research groups related to the development of oxyfuel combustion systems and for

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indirect-fired (closed loop) sCO2 power cycles [12-25]. In general, these studies have shown that despite its lower oxygen activity, pure CO2 at high temperature is considerably more corrosive to

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Fe-Cr steels than oxygen or air. The mechanism of this enhanced corrosion, whereby a slowgrowing Cr-rich oxide (henceforth referred to generally as “chromia”) scale is replaced by fast-

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growing Fe-rich oxides accompanied by internal carburization and significantly higher oxidation rates, has been described in detail [13, 26, 27]. The addition of secondary oxidants to the CO2 gas can significantly affect the oxidation behavior of the steel. This has been demonstrated in recent years by several research groups who

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are evaluating these materials for use in the aforementioned applications [12, 28-33]. A primary consideration is whether each impurity species serves to stabilize or disrupt the formation and growth of a protective chromia scale under the conditions of interest. In particular, the combustion impurities of H2O, O2, and SO2 have shown a complex pattern of behavior, where each has proven to be both beneficial and detrimental, depending on the precise exposure conditions (i.e., alloy

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composition, gas composition, and temperature). For example, the addition of H2O to CO2 is detrimental in that it can accelerate chromia growth [34], creates stresses in the chromia scale leading to local failure [35], and increases the amount of Cr in the steel that is required to form a chromia scale [12]. The last effect is associated with an increased extent of internal oxidation, which reduces the amount of Cr available for sustained chromia growth [12]. Alternatively, H2O additions can be beneficial by reducing the extent of internal carburization, particularly beneath

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Fe-rich oxides in the event of chromia failure [26, 34, 36]. This appears to be related to a

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competitive adsorption process, where H-containing species exclude CO2 and/or CO molecules from the oxide surface and/or internal surfaces such as oxide grain boundaries [37], which

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constitute the pathways for C ingress through the oxide scale [26, 28, 34]. The presence of O2 can

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likewise decrease the extent of internal carburization (and sulfidation when SO2 is present) by reducing the activity of carbon (and sulfur) in the gas, or can otherwise promote protective chromia

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formation [36]. However, O2 in combination with H2O accelerates chromia volatilization, which can in turn lead to chromia failure [34]. Finally, small amounts of SO2 can promote chromia

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formation and maintenance by reducing the amount of H- and C-containing species diffusing through the scale, by the competitive adsorption processes described above [30, 38-40]. This would reduce any negative impact these species have on the mechanical integrity of the chromia scale, as well as to reduce the extent of internal oxidation, internal carburization, and chromia

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volatilization. However, excess levels of SO2 can lead to formation of S compounds within or below the chromia scale, which can in turn cause failure [41]. As a result of these competing effects, the combination of SO2 and H2O has even been found to improve oxidation resistance compared to both pure CO2 and to CO2 containing only one of these impurities [41].

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Alloy composition also plays an important role in the ability of the steel to form and maintain a protective chromia scale in high temperature CO2 containing impurities. Aside from the obvious benefit of increased Cr content, minor additions of Mn and Si significantly influence the corrosion behavior. This was illustrated through a systematic investigation using model alloys of Fe-Cr and Fe-Cr-Ni doped with Mn and/or Si during exposure to CO2 with and without H2O and/or SO2 [38, 39, 42-49]. As with the effect of secondary oxidants, additions of Mn and Si

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produced both beneficial and detrimental effects, depending on the exposure conditions [14]. For

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example, in many of the alloys, Mn and Si both slowed chromia growth and reduced alloy carburization (thereby preventing scale failure) by forming layers of Mn-rich spinel and SiO2 that

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acted as diffusion barriers at the top and bottom of the chromia scale, respectively. Indeed, the

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beneficial effects of this silica layer for steels exposed to CO2 have long been known [50, 51]. While Mn improved performance of most of the above model alloys in CO2 containing H2O, the

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addition of SO2 caused increased oxidation in Fe-9Cr-2Mn (wt%) compared to the undoped alloy. The authors suggested that the high thermodynamic stability of Mn-sulfides, combined with the

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fast diffusion behavior of Mn cations within the chromia scale, potentially caused sulfides to form within the oxide at the early stages of reaction [38]. Likewise, while Si also generally improved oxidation performance, a high Si content (0.5 wt%) in combination with H2O caused growth stresses resulting in local failure of the chromia scale in austenitic Fe-Cr-Ni alloys [46]. Further,

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any amount of Si (0.1-0.5 wt%) in these alloys caused spallation during cooling, due to a combination of the above growth stresses and the thermal stresses that arose from the large difference in thermal expansion coefficients between the oxide and the austenitic alloy matrix [46]. Likewise, the addition of Si to both ferritic Fe-Cr and austenitic Fe-Cr-Ni alloys in the presence of

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H2O and SO2 caused oxide spallation on cooling due to the segregation of S to the interface of the alloy substrate and the SiO2 layer that formed at the base of the oxide [39]. The above work highlights the myriad of competing mechanisms in effect during high temperature oxidation of steels in impure CO2 that can both facilitate and impede protective behavior depending on the gas composition, alloy composition, and reaction temperature. Clearly, more work is needed to approach a mechanistic understanding of these processes over a wide range

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of the above experimental variables. Simultaneously, the extreme complexities encountered for

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even model alloys and gas compositions illustrate the need to evaluate the actual many-component commercial alloys at the precise exposure conditions of interest when considering these materials

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for new applications.

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In this study, we exposed several commercially available steels considered as candidates for construction of next-generation CO2-based power systems. Atmospheric pressure exposures at

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550 °C and a gas composition of 95% CO2, 4% H2O, 1% O2, with and without 0.1% SO2, were used to simulate the environment expected in intermediate-to-high temperature portions of the

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primary heat exchanger in a direct-fired sCO2 power cycle fueled by the combustion of natural gas and coal syngas, respectively [52]. 2. Experimental

Table 1 provides the compositions of the steels used in this study. The detailed

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experimental procedure was reported previously [53]. In short, exposure coupons of dimension 19 × 12.7 × 2 mm were machined from steel plates, which were purchased from the manufacturers in the service-ready condition. A 4.8 mm diameter hole was machined near the top middle of the coupon for hanging during exposure. The coupons were finished using 600 grit (CAMI designation) SiC paper, ultrasonically cleaned in acetone or alcohol, and weighed before exposure

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in a horizontal tube furnace to a gas composition of 95% CO2, 4% H2O, 1% O2, with and without 0.1% SO2 at 550 °C and 1 atm. The (pure) gases were mixed using individual mass flow controllers, then passed through a Pt-Rh catalyst mesh prior to reaching the samples, while the H2O was introduced by syringe pump as liquid water. The samples were cooled to room temperature and weighed every 500 h. The total time for the exposure done without SO2 was 2500 h. The final 500 h of the exposure done with SO2 was interrupted resulting in a total exposure time

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of 2250 h.

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After exposure, the samples were characterized using the following techniques. Surface and cross-sectional scanning electron microscopy (SEM) imaging was done using a FEI Inspect

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F50 SEM operating at 10-20 kV. Surface imaging was done using secondary electrons and cross-

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sectional imaging was done using backscattered electrons (BSE). The procedure for preparing sample cross-sections was described previously [53]. Cross-sections were imaged before and after

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etching using Murakami’s Reagent (ASTM standard etchant number 98) [54] to assess for carburization. Electron microprobe (EPMA) qualitative X-ray mapping and quantitative point

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analysis of the cross-sections were done using a JXA-8530FPlus HyperProbe. Scanning transmission electron microscopy (STEM) and associated energy-dispersive X-ray spectroscopy (EDS) was done using a FEI Titan 80-200 TEM equipped with a ChemiSTEM EDS detection system. A high-angle annular dark-field (HAADF) detector was used for imaging. The TEM

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samples were prepared by the focused ion beam (FIB) lift-out method using a FEI Helios Nanolab SEM.

3. Results

3.1 Mass change

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The mass change for all alloys exposed with and without SO2 for up to 2500 h are shown in Fig. 1. The steels that contained low (<9 wt%) Cr experienced high mass gains (7-8 mg/cm2) after 2500 h, with slightly less mass gain for Grade 91 compared to Grade 22. Comparison between Fig. 1a and 1b reveals effectively no difference between these two low-Cr steels for exposures done with and without 0.1% SO2. The steels that contained high (>17 wt%) Cr experienced very low mass gains (<0.1 mg/cm2) after 2500 h, on the order of what was detectable using the

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microbalance, for exposures done without SO2 (Fig. 1c). However, in contrast to the low-Cr steels,

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the presence of SO2 caused a significant increase in mass gain. While these mass gains are still relatively low (<1 mg/cm2), the results indicate a clear negative effect of SO2 for high-Cr steels.

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To shed light on the processes controlling this oxidation behavior, we selected several of the steels

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(indicated by dashed circles in Fig. 1) for characterization, presented below.

3.2 Microstructural characterization of Grade 91

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Surface and cross-sectional SEM images of Grade 91 exposed with and without SO2 are shown in Fig. 2. The image of the sample exposed without SO2 (Fig. 2a) reveals a relatively smooth surface interrupted by large regions of oxide spallation. Alternatively, the surface of the sample exposed with SO2 (Fig. 2b) is covered with small oxide nodules but shows little or no spallation.

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Cross-sectional images of the samples are shown in Fig. 2b and 2d. Both samples formed a duplex oxide scale consisting of an outer layer of Fe2O3/Fe3O4 and inner layer of Fe3-xCrxO4 mixed spinel oxide. This duplex oxide scale is consistent with extensive studies on the oxidation of 9-12Cr steels in pure high-temperature CO2 [15, 23]. Comparison between Fig. 2b and 2d reveals that a thicker oxide formed in the presence of SO2, and the difference in thickness is larger than would be

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expected based on the mass gains of the two samples (Fig. 1). An internal oxidation zone (IOZ) developed immediately below the oxide scale. The IOZ formed in the sample exposed with SO2 was both much thinner, and in some regions completely absent, compared to the sample exposed without SO2. Cross-sectional X-ray maps showing a typical region of the Grade 91 sample exposed with SO2 are provided in Fig. 3. The maps confirm the outer oxide layer is exclusively Fe-oxide, while

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the inner oxide layer also contains Cr. In addition, Fig. 3c suggests the IOZ is enriched in Cr, most

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likely as Cr-rich spinel oxide [55]. Figure 3d reveals some O enrichment near the top of the oxide scale, seen as darker regions in the BSE image (Fig. 3a). These regions are not correlated with the

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presence of any minor alloying elements (maps not shown), suggesting they are regions of Fe2O3

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located near the surface of the oxide. Figure 3e reveals a small amount of S located just below the surface of the oxide, and at the interface of the inner and outer oxide layers. In particular, the

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highest S content was observed as particles present at this interface, such as the region indicated in Fig. 3a. Analysis of this region at higher resolution is shown in Fig. 3f-j. The regions of highest

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S content correspond to particles immediately above the inner oxide layer. Quantitative EPMA analysis of one such particle (Fig. 3f) yielded a composition (at%) of 48.9% Fe, 43.8% S, 6.1% O, and other elements at trace levels. The size of this particle (≈2 µm) is beyond the spatial resolution of the technique, and therefore the surrounding oxide matrix contributes to the measured

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composition. Assuming the 6.1% O can be assigned to a surrounding oxide of pure Fe3O4, the particle itself would have a composition of 44.3% Fe and 43.8% S. This stoichiometry suggests particles of FeS have formed at the inner/outer oxide interface, consistent with previous observations [38, 55, 56].

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Grade 91 and similar steels carburize when exposed to high temperature CO2 [57]. To determine the extent of carburization, we etched the sample cross-sections and the results are shown in Fig. 4. Etching revealed a zone of carbides immediately below the oxide scale that extended >100 µm into the underlying alloy substrate (Fig. 4a-b). A higher magnification BSE image (Fig. 4c) illustrates the typical oxide/alloy interface, where carbides evident as dark contrast particles decorate the alloy. To quantify the extent of carburization, we acquired similar BSE

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images as a function of depth into the alloy at three random locations across each sample, then

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measured the fraction of carbides using ImageJ software. Figure 4d summarizes the results of this analysis. Both samples contained 9.0 area% of carbides at the surface of the alloy, immediately

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below the IOZ. Interestingly, the extent of carburization increases with depth for the sample

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exposed without SO2, reaching a maximum of 10.2 area% approximately 30 µm into the alloy. Alternatively, the sample exposed with SO2 shows a continuous decrease in the extent of

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carburization with depth into the alloy. A baseline level of metal carbides was seen for both samples, which were present prior to exposure and/or formed by thermal aging. This level was

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defined as the average carburization of the two samples at the middle of the metal section (940 µm below the oxide), indicated by the dashed line in Fig. 4d. Using this criterion, the carburization depth of the samples exposed without and with SO2 is approximately 150 and 200 µm, respectively. Based on these results and assuming carbides of Cr23C6 [15], the total mass of carbon

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uptake (determined by the area under the curve in Fig. 4d) for the samples exposed without and with SO2 was 0.40 and 0.22 mg/cm2, respectively. To determine the total expected mass gain of the Grade 91 samples, we also calculated

expected mass gains due to the external oxide scale (assuming pure Fe3O4) and internal oxide precipitates (assuming pure FeCr2O4), by measuring the area of these features at several regions

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across the samples shown in Fig. 2. The expected mass gain due to the scale was 5.9 and 7.5 mg/cm2, while that due to internal oxides was 0.9 and 0.3 mg/cm2, for the samples exposed without and with SO2, respectively. The sum of the calculated mass gains due to external scale, internal oxides, and internal carbides yielded a total estimated mass gain that was +0.4 mg/cm2 (+6%) and +1.0 mg/cm2 (+14%) relative to the measured mass gain (Fig. 1) for the samples exposed without and with SO2, respectively. The close agreement between the calculated and measured mass gains

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for the sample exposed without SO2, which is likely within the error of the measurement, suggests

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relatively little mass was lost due to oxide spallation, despite the surface observations (Fig. 2a). The slightly higher calculated mass gain for the sample exposed with SO2, with no evidence of

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spallation (Fig. 2c), suggests some local variation in the extent of oxidation across the sample

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surface that was not fully captured by the 2-dimensional cross-section.

3.3 Microstructural characterization of 347H

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Surface and cross-sectional SEM images of 347H exposed with and without SO2 are shown in Fig. 5. A thin Cr-rich oxide (chromia) scale with occasional small Fe-rich oxide nodules formed for the sample exposed without SO2 (Fig. 5a-b), confirmed by EDS measurements (not shown). In contrast, frequent large oxide nodules formed for the sample exposed with SO2 (Fig. 5d-e). In

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addition, small sulfide particles formed in the alloy substrate, as labeled in Fig. 5e. Qualitatively, these sulfide particles appeared more prevalent beneath the thick oxide nodules than beneath the thin chromia scales. After initial inspection, the cross-sections were etched and images acquired at the same region of the samples (Fig 5c, 5f). Comparison of Fig. 5b and 5c reveals a small amount of carburization of the alloy substrate for the sample exposed without SO2, to a depth of

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approximately 3 µm beneath the Cr-rich oxide scale. Alternatively, comparison of Fig. 5e and 5f reveals almost no carburization for the sample exposed with SO2, where most of the dark contrast particles near the surface of the alloy substrate are sulfides. Additional imaging of this etched sample (not shown) confirmed the existence of occasional carburization, but at even smaller amounts than the sample exposed without SO2. Cross-sectional X-ray maps showing elements of interest for a typical region of the 347H

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sample exposed with SO2 are provided in Fig. 6. The maps confirm that the thin oxide scale (left

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side of image) is Cr-rich, while the thick oxide nodule (right side of image) is Fe-rich. Remnants of a thin, broken chromia layer exist near the bottom of the Fe-rich nodule, flush with the still thin

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Cr-oxide scale near the left of the image. This suggests that a chromia scale formed initially in this

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region, then failed at some later time during the exposure. Further, since this layer serves as a marker of the original alloy surface, the thick oxide nodule constitutes both inward and outward

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oxide growth. Figure 6f shows that S exists to some extent throughout the entire oxide, as well as in the underlying alloy substrate. In particular, the highest concentration of S is observed within

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the inward growing region of the oxide nodule and in the alloy immediately below this region, whereas little S is observed in the vicinity of the thin Cr-rich oxide scale. Close inspection reveals that the S in the alloy is correlated primarily with Mn, while that in the inner oxide layer is correlated with Cr and Mn, as indicated with arrows in Fig. 6. Figure 6h also shows Nb present as

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particles throughout the alloy and inner oxide layer. The particles in the alloy correspond to Nbrich carbides, which were present prior to exposure. The size and distribution of Nb-rich particles within the inner oxide layer is consistent with that observed in the alloy, suggesting they are primarily the result of pre-existing carbides that were oxidized and incorporated into the oxide as it grew inwards. While it is difficult to exclude the possibility of Nb sulfides within the inner oxide

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layer at this resolution, TEM analysis shown below (Section 3.5) confirmed that Nb was present in these regions predominantly as an oxide. Fig. 6j also shows that a layer of silica has formed above the oxide nodule, as well as above the initially formed chromia scale present at the original alloy surface. This appears to be related, at least in part, to contamination from the mullite tube used during the exposure (similar analysis of a Si-free alloy exposed during the same test revealed the presence of small amounts of Si in the oxide scale). Finally, regions enriched in Ni appear in

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the alloy immediately below the oxide scale. We revisit this observation in Section 4.3.

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3.4 Microstructural characterization of 310S

We performed analogous characterization for 310S—another austenitic stainless steel that

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showed a less pronounced detrimental effect of SO2 (Fig. 1d). The surface and cross-sectional

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SEM results are summarized in Fig. 7. Comparison of Fig. 7a-b and Fig. 7d-e reveals largely similar results to 347H. That is, the presence of SO2 causes failure of the initially thin chromia

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scale, resulting in thicker Fe-rich oxide nodules. Likewise, sulfide particles appear for the sample exposed with SO2 (Fig. 7e), and these particles exist preferentially beneath the thicker Fe-oxide

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nodules. Comparison between Fig. 7d and Fig. 5d reveals that Fe-rich oxide nodules are somewhat less frequent for 310S relative to 347H, however some oxide spallation is also observed for 310S. This could be related to the relatively high Si content of 310S—excessive levels of Si have been linked to spallation of austenitic stainless steels during exposure in CO2 [51] and CO2-H2O gas

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mixtures [46], an effect which can be enhanced by the presence of SO2 [39]. To determine whether the lower measured mass gain relative to 347H (Fig. 1) could be attributed to spallation, the expected mass gain of 310S was calculated by measuring the oxide thickness at several regions across the sample. Assuming all mass gain was due to growth of the external oxide scale (i.e., the mass gain due to internal precipitates was negligible) and assuming a scale of pure Fe2O3 (the

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primary reaction product determined by X-ray diffraction—data not shown), the calculated mass gain was -0.01 mg/cm2 (-7%) relative to the measured mass gain. This very close agreement suggests that spallation contributed only minimally, and therefore the lower mass gain of 310S reflects less Fe-rich oxide growth compared to 347H, as discussed further in Section 4.4. The etched 310S sample exposed without SO2 (Fig. 7c) shows a similar depth of intragranular carburization compared to 347H (Fig. 5c). Carbides are also seen along grain boundaries in the

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alloy periodically across the sample surface, such as that shown in Fig. 7c. However, a similar size

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and distribution of carbides also existed along grain boundaries 1.0 mm below the oxide, at the middle of the metal section (not shown). Therefore, similar grain boundary carbides were present

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prior to exposure and/or formed by thermal aging, and we cannot determine to what extent

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intergranular carburization occurred during the reaction. The 310S sample exposed with SO2 (Fig. 7f) did show somewhat more intragranular carburization than the equivalent 347H sample (Fig.

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5f). The depth of carburization was similar compared to 347H and 310S exposed without SO2, however the carbide particles were smaller.

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Cross-sectional X-ray maps showing elements of interest for a typical region of the 310S sample exposed with SO2 are provided in Fig. 8. As with 347H, the thickest regions of the scale represent failure in protective chromia leading to growth of Fe-rich oxides. Likewise, S is present in small quantities throughout the oxide, and at relatively large quantities near the bottom of the

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oxide and in the underlying alloy, which is correlated primarily with Mn and Cr. Silicon is also seen near the surface of the oxide, related at least in part to contamination from the reaction tube as described above (note that the X-ray maps are individually normalized, i.e., qualitative). 3.5 TEM characterization of 347H

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The above results show that for high-Cr steels, SO2 in the reaction gas leads to formation of S-containing compounds and is associated with failure in protective chromia growth in favor of Fe-rich oxides and higher oxidation rates. To learn more about the mechanism behind this behavior, we performed TEM analysis on a region of 347H exposed with SO2 that was at a relatively early stage of nodule formation, as shown in Fig. 9a. The cross-sectional STEM image corresponding to the dashed line in Fig. 9a is shown in Fig. 9b, while X-ray (EDS) maps of the

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same region are shown in Fig. 9c-e. Figure 9d reveals a thin (<1 µm) layer of Cr-rich oxide

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immediately above the alloy, which represents the protective chromia scale formed initially during the exposure, while an Fe-rich oxide nodule has begun to grow above the chromia scale (Fig. 9c).

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To compare the oxide/alloy interface of the still-protective chromia scale with that which has

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experienced the onset of nodule formation, higher resolution analyses of the indicated regions in Fig. 9b are shown below.

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Figure 10 shows EDS maps for the thin oxide scale adjacent to the nodule. The maps confirm the scale is primarily chromia with minimal Fe and reveals some enrichment of Mn and

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Si at the bottom and top of the oxide, respectively. Figure 10f shows that a small amount of S is distributed throughout the oxide, and indeed quantitative analysis of these EDS maps indicated approximately 1 at% of S present in the chromia scale. Additional S is indicated near the oxide/alloy interface that is correlated with Mo and Nb and seen as bright contrast particles in the

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STEM image (Fig. 10a). However, careful examination of the X-ray spectra (Fig. S1 of the supplementary information) confirmed that this was an artifact of imperfect peak deconvolution, causing S (Kα = 2.31 keV) to be observed everywhere that Mo (Lα = 2.29 keV) or Nb (Lβ = 2.26 keV) were present. The only S truly present in the alloy in this region was in the form of a single Mn sulfide particle, as labeled in Fig. 10. The analysis in Fig. 10 also reveals a zone of

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microstructural transformations that extends ≈2 µm below the growing oxide scale. The origins of these transformations are discussed in section 4.3. Figure 11 shows similar EDS maps taken beneath the right side of the oxide nodule. The red dashed lines delineate the surface of the oxide, as well as the interface of the originally formed chromia scale and outward growing Fe-rich nodule. In contrast to the nodule-free chromia scale (Fig. 10), a significant quantity of Mn sulfide particles appears in the alloy immediately below the

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nodule, as indicated by the arrows in Fig. 11. In addition, Nb-rich particles have formed that span

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from the underlying alloy to above the Cr-rich oxide scale. While there appears to be some S correlated with Nb, we again confirmed this was purely an artifact of imperfect peak deconvolution

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of the X-ray spectra (Fig. S1 in the supplementary information). Thus, Nb within the scale was

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present exclusively as an oxide. In addition, a Nb-oxide particle is present immediately below the chromia scale. Close inspection shows that the bottom right of this particle contains Fe and no O,

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suggesting it was an Fe2Nb intermetallic particle that had begun to oxidize internally. The origin of these particles is discussed further in Section 4.3. While the participation of Nb introduces some

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complexity to the interpretation of these results, it remains clear that failure in protective chromia growth is associated with significant Mn sulfide formation in the underlying alloy. 4. Discussion

4.1 Thermodynamic considerations

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The experimental results show that oxide scales are the primary reaction product formed

during the exposures. In addition, C- and S-containing compounds also formed, to a varying extent, for all the steels investigated. In consideration of the thermodynamic driving forces for the formation of these compounds, Fig. 12 presents thermochemical diagrams calculated using the FactSage software package [58]. The plots show the stable compounds of individual alloying

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elements of interest in their pure form (i.e., Fe, Cr, etc.) as a function of the oxygen activity (pO2) and of the carbon (ac) or sulfur activity (pS2) at 550 °C. For clarity, only the most metal-rich compounds are shown for each element. The point corresponding to the equilibrium gas composition is indicated for each plot. Figure 12a shows that the oxides are stable relative to carbides for all metals in equilibrium with the gas. Alternatively, the equilibrium gas composition is near the transition in stability between the oxides and sulfates of Fe, Cr, Ni, and Mn. Assuming

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a perfectly equilibrated gas, stable compounds are the sulfates of Mn, Ni, and Cr, and the oxides

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of Fe. The cross-sectional X-ray maps of the stainless steels (Figs. 6 and 8) reveal a small amount of S at the surface of the oxides, which is correlated with Mn and Ni in addition to Cr. While these

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compounds were not identified directly, the results do suggest the possibility of small amounts of

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Mn and/or Ni sulfates at the outermost surface of the oxide. Alternatively, Grade 91, which contained less Mn and Ni and formed a much thicker oxide scale, showed no evidence of S

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compounds at the outermost oxide surface.

The compounds expected to form in contact with the gas can be predicted, at least in theory,

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based on the equilibrium gas composition as described above. However, this analysis provides no information on the compounds expected to form within and below the oxide scale, where the pO2 decreases as a function of depth. In particular, the existence of carbides and sulfides within the alloy implies that C- and S-containing species, respectively, have diffused through the growing

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oxide. Assuming local thermodynamic equilibrium, the pO2 in equilibrium with CO2 at any point within or below the oxide scale can be defined by the dissociation reaction: CO2(g) = CO(g) + ½O2(g)

(1)

while the carbon activity (ac) can be defined the tendency of CO produced by (1) to recombine via the Boudouard reaction:

17

2CO(g) = CO2(g) + C(s)

(2)

Likewise, recognizing that SO3 is a significant reaction product at these temperatures, the pO2 in equilibrium with pS2 can be similarly defined by the following reactions: SO2(g) + ½O2(g) = SO3(g)

(3)

SO2(g) = ½S2(g) + O2(g)

(4)

While a more rigorous description of this process can be found in the reference [59], the important

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point is that local thermodynamic equilibrium dictates that ac and pS2 increase as pO2 decreases

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with depth into and below the oxide scale. Thus, successful diffusion of C- and S-containing species through the oxide produces elevated levels of ac and pS2.

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The solid black lines in Fig. 12 represent the hypothetical scenario where the reaction gas

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successfully diffuses through the oxide scale and achieves local equilibrium with the pO2 that exists at every point. Thus, the surface of the oxide is defined by the equilibrium gas composition,

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while moving left along this line defines the hypothetically stable compounds for increasing depth into and below the oxide scale. Figure 12a shows that at sufficiently low pO2 values, the stable

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compound changes from oxide to carbide for all elements of interest. More specifically, the vertical lines labeled Fe(s) and Cr(s) represent the pO2 value defined by the metal/oxide interfaces of Fe/Fe3O4 (FeO is not stable at 550 °C) and Cr/Cr2O3, respectively. As the decreasing pO2 approaches and drops below these values, the stable compound changes from oxide to carbide for

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both Fe and Cr. This is consistent with the experimental results, where metal carbides precipitate in the alloy beneath both Fe-rich (Fig. 4) and Cr-rich (Figs. 5 and 7) oxide scales. Similar analysis in Fig. 12b shows that some metal sulfates are stable at the surface of the

oxide, however the stable compound transitions from sulfate to oxide for relatively small decrease in pO2, suggesting that sulfate would only be expected near the outermost surface of the scale.

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Moving left along the line in Fig. 12b shows that the oxides of all metals are stable for approximately 10 orders of magnitude decrease in pO2, below which the stable compound of all metals change from oxide to sulfide. This transition occurs at pO2 values significantly higher than those defined by the Fe3O4/Fe and Cr2O3/Cr equilibrium values. Therefore, the existence of sulfides within the oxide scale is predicted, assuming the successful transport of S-containing species through the oxide. This is consistent with the observation of FeS particles throughout the

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Fe-rich oxide formed on Grade 91 (Fig. 3). In addition, the observation of S throughout both the

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thin Cr-rich oxides and the thick Fe-rich oxide nodules for the high-Cr steels (Fig. 6 and Fig. 8) suggests that sulfides of Mn and/or Cr have formed to some extent throughout most of the oxide

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scale. Figure 12b shows that the relative stability of metal sulfide compounds follows Mn > Nb >

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Cr > Mo > Fe > Ni. This is consistent with the observation that Mn sulfide particles were the only S compounds that formed internally beneath the Cr-rich oxides present on the high-Cr steels, where

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the pS2 is lowest. Despite their relatively high stability, little or no Nb sulfides formed. Instead, Nb oxides formed in regions of chromia failure, either by outward diffusion through the chromia

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(Fig. 11) or by the oxidation of Nb-rich carbides as they were incorporated into the inward growing portion of the scale (Fig. 6). Besides Mn, Cr was the only other element that was significantly correlated with S for the high-Cr steels—primarily within the inward growing portion of the scale (Fig. 6 and Fig. 8). However, no Cr sulfides were found beneath the oxide, both for the still-

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protective chromia scale (Fig. 10) and for the early stages of nodule growth (Fig. 11). Thus, Cr sulfides likely formed within the oxide scale itself, rather than formation in the alloy substrate with subsequent incorporation into the inward growing scale. 4.2 Competitive adsorption of reactive molecules

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Many experimental observations have led to the proposal that different molecular species present in multi-oxidant environments compete for adsorption sites on oxide surfaces during high temperature oxidation [59]. This can, for example, change the oxidant activity at the outermost surface of the oxide to be different from that of the gas. More importantly, the most likely mechanism for C and S transport through the oxide, which is required to achieve the elevated activities necessary to form carbides and sulfides as described in Section 4.1, appears to be the

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molecular transport of these species along oxide grain boundaries or similar internal surfaces [37].

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Therefore, the preferential adsorption of certain of these molecules at these internal surfaces can dramatically affect the oxidation behavior.

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Grade 91, which formed thick Fe-rich oxides, underwent spallation and formed a relatively

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deep internal oxidation zone during exposure without SO2 (Fig. 2). Both of these are related in part to the presence of H2O in the reaction gas, which is known to cause stresses in the oxide and to

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increase the extent of internal oxidation [12]. However, in the presence of 0.1% SO2, both spallation and internal oxidation are reduced (Fig. 2), while Fe-rich sulfides are formed in the scale

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(Fig. 3). This suggests that S-containing species have successfully displaced H-containing species diffusing inward through the oxide during the exposure. Similarly, the extent of internal carburization of Grade 91 (Fig. 4) was reduced in SO2, suggesting also the partial displacement of C-containing species leading to a lower ac at the oxide/alloy interface [30, 38, 39]. Finally, the

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more extensive internal oxidation of Cr for exposure without SO2 led to Cr depletion near the alloy surface, causing the maximum carbide volume fraction to occur ≈30 µm below the oxide/alloy interface. Alternatively, the reduction in internal oxidation for the exposure with SO2 led to minimal Cr depletion, and therefore a maximum carbide volume fraction at the oxide/alloy interface where ac was highest.

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For the case of the high-Cr steels, exposure without SO2 resulted in thin chromia scales and a small amount of internal carburization. As with Grade 91, the presence of SO2 reduced the extent of internal carburization by displacing inward diffusing C-containing species. However, this resulted also in the formation of sulfides within and below the chromia scale and ultimately failure in protective behavior, as discussed further in Section 4.4. 4.3 Evolving microstructure of the alloy surface

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TEM analysis of the thin chromia scale (Fig. 10) reveals many structural transformations

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near the surface of the alloy that are associated with the oxidation process. We now consider these in more detail, as shown in Fig. 13. First, voids formed at the oxide/alloy interface. These voids

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form by the injection of vacancies into the alloy substrate as metal atoms are ionized and diffuse

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outward during oxide growth [60].

Second, several bright contrast particles are seen near the oxide/alloy interface in Fig. 13a

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that are rich in Nb and/or Mo (Fig. 10g-h). These regions were not associated with increased levels of C, suggesting the particles were not carbides. Instead, they appear to be intermetallic

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Fe2(Nb,Mo) Laves phase particles. The precipitation of these phases at the oxide/alloy interface is consistent with previous investigations of austenitic stainless steels exposed to high temperature CO2 [61]. The driving force for formation of these compounds arises from the decreased Cr activity at the alloy surface, which accompanies growth of the chromia scale. This induces the uphill

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diffusion of Nb toward the oxide/alloy interface, with subsequent precipitation of the Nb-rich intermetallic phases [62]. Third, a zone consisting of several recrystallized grains exists within the top ≈2 µm of the

alloy surface. These likely formed from stresses incurred at the alloy surface during grinding, with subsequent recrystallization and/or grain growth by thermal aging during the exposure [63]. In

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addition, the diffusion processes associated with oxidation might also contribute to surface recrystallization [64]. Regardless of its origin, this recrystallization zone plays an important role during the oxidation process. From the EDS maps, the composition of primary alloying elements Fe, Ni, and Cr clearly vary among the different grains and relative to the underlying substrate. Figure 13g compares the compositions of 6 recrystallized grains (labeled in Fig. 13a) as well as the substrate immediately below these grains. The substrate composition is consistent with that of

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the bulk alloy (Table 1), suggesting the supply of Cr supporting formation/growth of the chromia

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scale came exclusively from the recrystallized grains. In addition, Fig. 13g reveals a bimodal distribution for Ni content among the recrystallized grains, where they contain either 3.1-3.6 wt%

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Ni or 17.3-18.5 wt% Ni. The substrate (and Ni-rich grains) exhibit an austenitic FCC structure, as

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shown in Fig. 13b. Alternatively, the Ni-poor grains have undergone phase transformation, resulting in BCC and/or BCT ferritic/martensitic structure. Further, these Ni-poor grains also

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contain less Cr (4.5-5.0 wt% Cr) compared to the Ni-rich grains (5.9-7.4 wt% Cr). This suggests that more Cr has diffused from these grains to support growth of the chromia scale, consistent with

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faster bulk diffusion in BCC/BCT compared to FCC matrices. In summary, structural transformations occur at the oxide/alloy interface during oxidation including void formation, secondary phase precipitation, and recrystallization. The surfaces/boundaries created from these processes provide fast diffusion paths for metal, which can

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result in profoundly different surface oxidation reactions than if controlled by bulk metal diffusion alone, especially at this relatively low temperature. For example, it is interesting to note that while a recrystallization zone is present beneath the thin chromia scale on either side of the oxide nodule, no such zone exists beneath the nodule itself (Fig. 9b). If local variations across the alloy surface result in the absence of this zone, these regions would presumably be more susceptible to Fe-rich

22

oxide growth, since they lack the fast diffusion paths for Cr to repair the chromia scale in the event of initial failure. These results highlight the importance of the structural damage induced by surface finishing (e.g., grinding with 600 grit SiC paper) for enabling thin Cr-rich oxide scales to form and grow on steels exposed to high temperature CO2 [63]. 4.4 Mechanism of chromia failure Figure 14a plots the mass gain versus Cr content of all tested steels, which illustrates that

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a transition from thick Fe-rich oxides (high mass gains) to thin Cr-rich oxides (low mass gains)

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occurred for an alloy Cr content between 9 and 17 wt%. Further, Fig. 14b shows that all alloys with enough Cr (>17 wt%) to form a protective chromia scale resulted in similar and low oxidation

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rates for exposures done without SO2. Alternatively, the addition of SO2 caused periodic failure in

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protective chromia growth, leading to frequent Fe-rich oxide nodules and higher oxidation rates for these high-Cr steels. Because the increased mass gain was largely attributed to growth of Fe-

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rich oxide nodules, the addition of SO2 had little effect on the low-Cr steels, which had already evolved to steady-state growth of thick Fe-rich oxide scales. For the high-Cr steels exposed with

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SO2, Fig. 14b shows that the alloy continues to benefit from an increase in Cr content beyond 17 wt%, up to and including the highest Cr-containing steel (≈27 wt%). This reflects that in the event of local chromia failure, alloys containing the highest Cr are most capable of re-passivating, resulting in less Fe-rich oxide growth. Indeed, the formation of a Cr-rich oxide “healing layer” at

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the base of the scale, after initial failure, is a known phenomenon for steels exposed to high temperature CO2 [65]. A clear example of this behavior is seen in Fig. 6. After initial scale failure resulted in Fe-rich oxide growth, a new Cr-rich oxide layer formed at the base of the scale (Fig. 6c). Diffusion through this newly formed Cr-rich oxide layer once again becomes rate limiting, slowing subsequent oxide growth. This sequence of behavior (Cr-rich oxide formation, failure,

23

and recovery) is reflected by the “S-shaped” curves for the mass gain of the samples shown in Fig. 1d. The precise mechanism by which SO2 causes failure in chromia growth is of obvious interest. Figures 10 and 11 reveal that only the sulfides of Mn form in the alloy beneath the chromia scale and during the early stages of nodule formation, and they are significantly more prevalent in the latter case. Therefore, the possibility of Cr sulfides forming in the substrate, thereby reducing

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Cr activity in the alloy and resulting in chemical failure of the chromia scale, does not appear to

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be a plausible explanation. The formation of a layer of Mn-rich oxide at the surface of the chromia scale has been linked to improved oxidation resistance in high temperature CO2 [14]. In our study,

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a layer of Mn-rich oxide formed instead at the bottom of the chromia scale (Fig. 10i). Still, similar

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Mn-rich oxides appeared at the bottom of the chromia scale present beneath the nodule (Fig. 11i), suggesting depletion of Mn in the substrate is also not responsible for chromia failure. Instead, the

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most likely explanation appears to be related to the presence of S within the chromia scale itself. This could lead to formation of sulfides within the chromia, which could deteriorate the mechanical

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integrity of the scale [38]. In addition, the formation of Cr-sulfides at the oxide grain boundaries would provide fast diffusion paths for Cr cations, thereby accelerating chromia growth [41]. This is consistent with the observation of Cr and Mn sulfides in the oxide for the case of the later stages of nodule growth (Fig. 6 and Fig. 8). Finally, the formation of sulfides and/or the preferential

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adsorption at oxide grain boundaries by S-containing molecules could also prevent the healing, via grain growth, of micro-cracks that form in the chromia scale. In short, the increased sulfide formation seen beneath the regions of chromia failure occurs due to a local increase in the

24

permeability of S in these regions, and is therefore associated with, but not the cause of, local scale failure. Upon local failure, Fe cations diffuse rapidly outward forming Fe-rich oxides above the chromia scale. In addition, Fig. 11 shows that Nb cations also diffuse outward, forming Nb-rich oxides within and above the chromia scale. The supply of Nb to the oxide/alloy interface to support this growth is provided by its uphill diffusion, as described in Section 4.3. Niobium oxide is more

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stable than Cr oxide (Fig. 12), and therefore the existence of Nb-rich oxides within and above the

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chromia scale is inconsistent with local thermodynamic equilibrium. This reflects the transient nature of the nodule, where the oxide has yet to reach its steady state morphology.

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We note that previous results of high-Cr steels 304H [30] and Fe-Cr-(Ni, Mn, Si) model

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alloys [38, 39, 41] exposed to similar CO2-H2O-SO2 gas mixtures at a higher temperature (650 °C) did not exhibit the same negative effects of SO2. Instead, SO2 reduced the corrosion rate [30] or

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even promoted full transition to thin chromia scales [38, 39, 41] compared to the SO2-free gas. The beneficial effect of SO2 is thought to be related to preferential adsorption of S- over C- and H-

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containing species on internal surfaces of the oxide (as discussed in Section 4.2), thereby reducing their negative impact on the mechanical integrity of the scale, and/or by reducing the extent to which they cause internal carburization and internal oxidation. Preferential adsorption of Scontaining species, and therefore its beneficial effect based on the above mechanism, would be

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expected to increase with decreasing temperature [40]. This is contrary to the results shown here at 550 °C. Therefore, it is possible that a certain amount of preferential adsorption is helpful, while excessive amounts become detrimental, causing SO2 impurities to be beneficial only in some intermediate temperature (or concentration) range. Recently, other authors proposed a similar explanation for why CO2 containing both H2O and SO2 resulted in successful chromia

25

formation/growth in Fe-20Cr model alloys, while equivalent exposures in pure CO2, CO2-H2O, and CO2-SO2 all resulted in chromia failure [41]. In this case, the amount of S in the oxide was controlled to be within the “beneficial range” by the addition of H2O, via the competitive adsorption processes described in Section 4.2. Alternatively, it is possible that SO3 is the reactive species responsible for the negative effects associated with the SO2-containing gas, since significantly higher quantities of SO3 (via Equation 3) exist at lower temperatures. Specifically,

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the ratio of SO3/SO2 in our reaction gas is 2.2 at 550 °C, whereas it would be 0.5 at 650 °C. Lower

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temperatures also result in less solid-state diffusion, both in the alloy and the oxide. This could make it more difficult to reform a chromia scale in the event of failure (due to slow Cr diffusion

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in the alloy), as well as reduce the ability of the chromia to heal from micro-cracks (in the form of

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oxide grain growth). Finally, metal sulfates are stable near the surface of the oxide (Fig. 12) at 550 °C, while they are not stable at higher temperatures. Therefore, a transition in the stable compound

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in equilibrium with the gas from metal oxides to metal sulfates, particularly for Cr, could be responsible for the failure in protective behavior at lower temperatures. However, chromia scales

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still persisted on at least portions of the sample surface, suggesting that this is unlikely to be the only mechanism in effect.

5. Summary and Conclusions

Six steels with a wide range of Cr content (2-27 wt%) were exposed to CO2-rich gas

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containing impurities of H2O, O2, with or without SO2, at 550 °C to simulate environments expected in the intermediate-to-high temperature portions of a direct-fired sCO2 based power cycle. Steels containing ≤9 wt% Cr formed thick duplex Fe-rich oxide scales in all cases. The presence of SO2 led to minor sulfide formation within the scale for these low-Cr steels but did not significantly affect the oxidation rate. Instead, SO2 reduced the extent of spallation, internal

26

oxidation, and internal carburization that are generally associated with H2O and CO2, respectively. Steels containing ≥17 wt% Cr formed thin Cr-rich oxide scales without SO2, whereas the presence of SO2 caused scale failure leading to frequent Fe-rich oxide nodules, and spallation for steels containing high levels of Si. This failure was associated with local increases in permeability of S through the oxide, which ultimately disrupted the mechanical integrity of the scale and/or accelerated the chromia growth rate. Thus, the presence of SO2 at these conditions is detrimental

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for the case of thin protective chromia scales but is beneficial after the steels have already evolved

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to a uniformly covered surface of Fe-rich oxide, by reducing the effects of secondary degradation modes.

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The results suggest that caution is required when considering stainless steels for use in

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applications where SO2 impurities accompany the CO2-rich gas. Further, comparison with previous investigations suggests that the effects of SO2 are especially sensitive to exposure

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conditions and can become more severe at lower temperatures. This emphasizes the need to evaluate alloys at the range of temperatures expected for a given application, which can span

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several hundred degrees Celsius. Finally, the results of this study highlight the many competing mechanisms controlling corrosion of steels in high temperature impure CO2 environments. Additional research is required to unravel these processes, to improve fundamental understanding

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and enable realistic lifetime predictions.

Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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☐The authors declare the following financial interests/personal relationships which may be considered

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as potential competing interests:

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Acknowledgements This work was performed in support of the U.S. Department of Energy’s Fossil Energy Crosscutting Technology Research Program. The Research was executed through the NETL Research and Innovation Center’s Advanced Alloy Development Field Work Proposal. Research performed by Leidos Research Support Team staff was conducted under the RSS contract

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sections and Keith Collins (NETL) for assistance with EPMA analysis.

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89243318CFE000003. We thank Christopher McKaig (NETL) for preparation of the sample cross-

Disclaimer

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This work was funded by the Department of Energy, National Energy Technology

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Laboratory, an agency of the United States Government, through a support contract with Leidos Research Support Team (LRST). Neither the United States Government nor any agency thereof,

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nor any of their employees, nor LRST, nor any of their employees, makes any warranty, expressed or implied, or assumes any legal liability or responsibility for the accuracy, completeness, or

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usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights. Reference herein to any specific commercial product, process, or service by trade name, trademark, manufacturer, or otherwise, does not necessarily constitute or imply its endorsement, recommendation, or favoring by the United States

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Government or any agency thereof. The views and opinions of authors expressed herein do not necessarily state or reflect those of the United States Government or any agency thereof.

Data Availability: All relevant data that support this study are available from the corresponding author upon request.

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Fig. 1 Mass change of steels exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2 and the same gas containing 0.1% SO2 for up to 2500 h. The plots are separated into (a-b) Low-Cr steels (c-d) High-Cr steels. Characterization results of the circled data points are presented throughout this paper.

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Fig. 2 Surface and cross-sectional SEM images of Grade 91 steel exposed at 550 °C and 1 atm to (a-b) 95% CO2, 4% H2O, 1% O2 (2500 h exposure) and (c-d) the same gas containing 0.1% SO2 (2250 h exposure).

Fig. 3 Qualitative X-ray mapping of Grade 91 steel exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a-e) low magnification BSE image and associated maps of Fe, Cr, O, S (f-j) higher magnification BSE image and associated maps of the region indicated in (a). Quantitative EPMA point analysis (at%) of a S-containing region is shown in (f).

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Fig. 4 Cross-sectional SEM images of etched Grade 91 samples exposed at 550 °C and 1 atm to (a) 95% CO2, 4% H2O, 1% O2 (2500 h exposure) and (b) the same gas containing 0.1% SO2 (2250 h exposure). (c) high magnification image typical of the oxide/alloy interface after etching (d) carbide area% measured from images such as (c) as a function of depth into the alloy substrate at three different locations across the sample. The error bars represent one standard deviation.

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Fig. 5 Surface and Cross-sectional SEM images of 347H exposed at 550 °C and 1 atm to (a-c) 95% CO2, 4% H2O, 1% O2 (2500 h exposure) and (d-f) the same gas containing 0.1% SO2 (2250 h exposure). The same regions of the sample cross-sections were used for imaging (b, e) before and (c, f) after etching.

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Fig. 6 Qualitative X-ray mapping of 347H exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a) BSE image and (b-j) associated maps showing elements of interest. Reference arrows are shown at the same locations on (a), (c), (f), and (i).

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Fig. 7 Surface and Cross-sectional SEM images of 310S exposed at 550 °C and 1 atm to (a-c) 95% CO2, 4% H2O, 1% O2 (2500 h exposure) and (d-f) the same gas containing 0.1% SO2 (2250 h exposure). The same regions of the sample cross-sections were used for imaging (b, e) before and (c, f) after etching.

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Fig. 8 Qualitative X-ray mapping of 310S exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a) BSE image and (b-h) associated maps showing elements of interest. Reference arrows are shown at the same locations on (a), (c), (f), and (g).

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Fig. 9 TEM analysis of an oxide nodule formed on 347H exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a) Surface SEM image showing the region used to prepare the TEM sample. (b) Cross-sectional STEM image of the oxide nodule in (a), with the surface of the oxide indicated by the red dashed line. (c-e) EDS maps of the region shown in (b). Higher magnification EDS maps of the regions indicated in (b) are shown in Figs. 10 and 11.

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Fig. 10 TEM analysis of the thin oxide layer adjacent to the oxide nodule shown in Fig. 9 for 347H exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a) STEM image and (b-j) corresponding EDS maps showing elements of interest. The arrows in (a), (f), and (i) point to an internal Mn-sulfide particle.

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Fig. 11 TEM analysis of the bottom of the oxide nodule shown in Fig. 9 for 347H exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a) STEM image and (b-j) corresponding EDS maps showing elements of interest. Mn sulfide particles are indicated with arrows in (a), (f), and (i).

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Fig. 12 Thermochemical diagrams calculated using FactSage software [58] for 550 °C and 1 atm showing pure metals (M) of interest in (a) the M-C-O system and (b) the M-S-O system. For clarity, only the most metal-rich compounds are shown.

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Fig. 13 TEM analysis of the structural transformation zone shown in Fig. 10 for 347H exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2, 0.1% SO2 for 2250 h. (a) STEM image (b) SAED pattern of substrate below recrystallization zone (c) SAED pattern of grain 3 (d-f) EDS maps of region in (a) showing elements of interest (g) EDS composition measurements (wt%) of the recrystallized grains labeled in (a).

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of ro -p re lP ur na Jo Fig. 14 Mass change at 2000 h exposure time versus chromium content for steels exposed at 550 °C and 1 atm to 95% CO2, 4% H2O, 1% O2 and the same gas containing 0.1% SO2 showing (a) all steels (b) expanded region showing only high-Cr steels. 47

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Table 1 Composition of the alloy test materials provided by the manufacturer (wt%). *Compositions for 304H and E-Brite were measured by wavelength dispersive X-ray fluorescence and the listed C contents are nominal values. Fe

Ni

Cr

Mo

Mn

Nb

Si

C

Other

Grade 22

95.52

0.15

2.29

0.94

0.52

-

0.21

0.13

0.03 Al, 0.17 Cu

Grade 91

89.31

0.09

8.37

0.90

0.45

0.07

0.33

0.09

0.01 Al, 0.22 V, 0.09 Cu

347H

70.13

9.01

17.34

0.37

1.86

0.52

0.31

0.05

0.14 Co

304H*

70.63

8.27

18.74

0.12

1.08

0.01

0.44

0.07*

0.01 Al, 0.22 Co, 0.01 W, 0.27 Cu

310S

53.49

19.10

25.04

0.09

1.39

0.01

0.39

0.04

0.02 Al, 0.17 Co, 0.18 Cu

E-Brite*

71.64

0.21

26.50

1.00

0.04

0.12

0.25

0.01*

0.10 Al, 0.02 Co, 0.07 V, 0.01 Cu

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Alloy

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