# The Third Element Effect during Aqueous Corrosion of FeCrAl Alloys: Effect of Al Content

Catherine S Lynch, Debashish Sur, Peter F Connors, Benjamin W Y Redemann, Junsoo Han, Elizabeth J Opila, John R Scully

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Catherine S Lynch, Debashish Sur, Peter F Connors, Benjamin W Y Redemann, Junsoo Han, et al. The Third Element Effect during Aqueous Corrosion of FeCrAl Alloys: Effect of Al Content. Journal of The Electrochemical Society, 2025, 172, 10.1149/1945-7111/ae20c9. hal-05489494

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# The Third Element Effect during Aqueous Corrosion of FeCrAl Alloys: Effect of Al Content

To cite this article: Catherine S. Lynch et al 2025 *J. Electrochem. Soc.* **172** 121501

View the [article online](#) for updates and enhancements.

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# The Third Element Effect during Aqueous Corrosion of FeCrAl Alloys: Effect of Al Content

Catherine S. Lynch,<fnref n="1" />,<fnref n="2" />,⁷ ID Debashish Sur,<fnref n="1" />,<fnref n="2" /> ID Peter F. Connors,<fnref n="1" />,<fnref n="2" /> ID Benjamin W. Y. Redemann,<fnref n="3" /> ID Junsoo Han,<fnref n="4" /> ID Elizabeth J. Opila,<fnref n="1" />,<fnref n="2" />,* ID and John R. Scully<fnref n="1" />,<fnref n="2" />,** ID

The *third element effect* draws on the synergy of some ternary systems. FeCrAl is one such alloy system that has shown superior high temperature oxidation resistance. However, the third element effect is understudied in aqueous corrosion. This study systematically explored the passivation behavior 4–16 at.% Al additions to single phase solid solution FeCr₁₀ alloys and compared their performance to Fe-Cr binary alloys. Additions of Al improved repassivation and long-term passivity of Fe<sub>90-x</sub>Cr<sub>10</sub>Al<sub>x</sub> alloys in deaerated acidified 0.1 M Na₂SO₄ or 0.1 M H₂SO₄. Characterization via *ex-situ* X-ray photoelectron spectroscopy and *in-operando* atomic emission spectro-electrochemistry explored compositions of the films as well as the specific elemental dissolution kinetics into the electrolyte. Oxides or hydroxides consisting of Cr₂O₃, Cr(OH)₃, and Al₂O₃ improved passivation of Fe<sub>90-x</sub>Cr<sub>10</sub>Al<sub>x</sub> alloys, relative to their ferrous and ferric oxides. The FeCr<sub>10</sub>Al<sub>16</sub> alloy composition behaved similarly to FeCr₁₆ when considering passivation performance. Even low concentrations of Al resulted in significant performance improvements in the passivity of FeCr₁₀ compared to binary Fe-Cr alloys with 16 at.% Cr. Discussion of the role of Al in the context of producing a third element effect, enhancing Cr (III) passivation, and as a secondary passivator itself is considered.

© 2025 The Author(s). Published on behalf of The Electrochemical Society by IOP Publishing Limited. This is an open access article distributed under the terms of the Creative Commons Attribution 4.0 License (CC BY, https://creativecommons.org/licenses/by/4.0/), which permits unrestricted reuse of the work in any medium, provided the original work is properly cited. [DOI: 10.1149/1945-7111/ae20c9]

<figure><img src="image_5.png" /></figure>

Manuscript submitted September 29, 2025; revised manuscript received November 16, 2025. Published December 5, 2025.

Supplementary material for this article is available online

In binary alloys, there typically exists a minimum alloying content (or critical threshold) of the primary passivating element that is required to achieve excellent passivation both at elevated as well as room temperature. This can be circumvented by adding a so-called *third element*. FeCrAl is one of the alloy systems in which the “third element effect” (TEE) is observed toward alloy passivity with stabilization of continuous protective α-Al₂O₃-rich scales upon Cr addition during exposure to oxygen containing environments at elevated temperatures.<fnref n="1" />⁻⁶ In general, this effect occurs when a *third element* (in this case, Cr) of intermediate thermodynamic oxide stability (Cr₂O₃) compared to that of the other first (Fe) and second (Al) metallic components is added to its binary alloy, to improve its oxidation resistance by stabilizing its second element’s oxide (α-Al₂O₃). The oxide of the first element is typically non-protective, while that of the second element serves as the primary passivator, as it forms a continuous protective oxide scale.⁶⁻⁸ Cr additions have been shown to reduce the Al threshold required from > 16 at.% in Fe to 6 at.% when > 20 at.% Cr is alloyed. Several works on commercial Kanthal type alloys⁹⁻¹¹ have demonstrated the significant effects of Cr in promoting α-Al₂O₃ formation and protection during exposure to dry as well as wet oxygen containing environments at elevated temperatures, but a much more limited understanding of FeCrAl passivation in ambient temperature aqueous corrosion exists.

The FeCrAl ternary alloy system is under consideration as an accident-tolerant fuel (ATF) cladding in light water reactors to replace Zircaloy. As shown by the Fukushima incident in 2012, there is a need for alloys designed for loss of coolant accidents,¹²,¹³ as Zircaloy can corrode in high temperature steam.¹⁴ At lower temperatures, it is more desirable to form Cr₂O₃-rich scales rather than the less protective Θ-Al₂O₃ scales. Interestingly, FeCrAl alloys

provide a stable and switchable dominant oxide protection strategy in both aqueous and high temperature steam environments, making them an attractive cladding substitute.¹⁵ These alloys have been shown to produce a thin and protective scale that can withstand elevated temperatures and submerged environments for prolonged periods of time.¹⁵,¹⁶

The research areas of aqueous corrosion and high temperature oxidation have long been separate fields with very minimal overlap. This separation, despite both fields concerning metal oxidation, is largely due to a difference in the thickness of oxide scales. High temperature scales are easily grown to the order of micrometers, whereas aqueous passive films are typically below 10 nm.¹⁷⁻¹⁹ This difference in length scale necessitates the use of different characterization methods for oxide films formed through high temperature oxidation and room temperature aqueous corrosion. Additionally, due to the nature of aqueous electrochemistry, hydroxides not only form but can also remain as the predominant oxides, with relevant properties to consider. Typically, high-temperature hydroxides will volatilize off of the surface rather than grow an adherent and protective layer.²⁰⁻²² Therefore, key differences in the mechanisms and properties of the TEE in the ambient temperature aqueous environment versus high temperature oxidation are a key knowledge gap in these converging fields. This work reports the properties of films formed by FeCrAl alloys in acidified sulfate solutions, aiming to identify the mechanisms behind the TEE in this compositional space systematically.

In aqueous solutions, primary passivating elements have a critical concentration threshold to overcome to form a complete and protective oxide passive film. For a Fe-Cr binary alloy, it has been observed that 12 at.% is the minimum amount of Cr in the bulk of the alloy necessary to form a coherent, passive film in sulfuric acid.²³,²⁴ That being said, this study chose the alloy to have 10 at.% to stay below this critical threshold to accentuate the effectiveness of the TEE. For binary Fe-Al alloys, the passivation is much weaker in acidified sulfate systems compared to Cr and has a much higher threshold of 18 at.% of Al.²⁵,²⁶

<footnote n="1">¹Center for Electrochemical Science and Engineering, University of Virginia, Charlottesville, Virginia 22903, United States of America</footnote>

<footnote n="2">²Department of Materials Science and Engineering, University of Virginia, Charlottesville, Virginia 22903, United States of America</footnote>

<footnote n="3">³Department of Materials Science and Engineering, Johns Hopkins University, Baltimore, Maryland 21218, United States of America</footnote>

<footnote n="4">⁴Laboratoire Interfaces et Systèmes Electrochimiques (LISE), Sorbonne Université, Paris 75005, France</footnote>

<footnote>*Electrochemical Society Member.</footnote>

<footnote>**Electrochemical Society Fellow.</footnote>

<footnote>E-mail: kte8kk@virginia.edu</footnote>
Recently, TEE type behavior has been reported toward aqueous passivity in ternary CuAlSn²⁷ and NiCrMo,²⁸ quaternary FeCrAlTi,²⁹ and multi-principal element alloys (MPEAs)³⁰⁻³² containing Al, Si, or Ti. This study was motivated to consider specific roles of elements observed in MPEAs to aid in alloy design. MPEAs have very complex interactions that can result in unexpected property changes; therefore, by studying ternary alloys, we can investigate interrelations between elemental properties without being muddled by the increased complexity of MPEAs.³³⁻³⁷ Systematic studies of elemental interactions of properties such as corrosion resistance are more valuable than traditional “cook and look” studies because the inter-dependencies uncovered can be used to further refine the modern alloy design process.

Overall, this study demonstrates the potency of the TEE as operative in FeCrAl with suboptimal chromium concentrations (below 10 at.%). Comparisons were made to binary Fe-Cr and Fe-Al alloys such as FeCr₁₆ and FeAl₁₆. Evidence that the TEE promotes primary passivation by minimizing Fe matrix oxidation, creating a scaffold for the primary or dominant passivator, as well as increasing the thermodynamic activity of the primary passivator, is shown. Therefore, this study elaborates on the TEE in ambient temperature aqueous corrosion in acidified sulfate solutions to understand the scientific mechanisms that can be used to design future alloys, from MPEAs to ATF cladding.

## Methods

Alloy ingots of Fe-Cr-Al of compositions listed in Table I were made to be 2.5 cm long with a 1 cm diameter via arc melting using pure Fe (purity 99.5%), Cr (99.99%), and Al (99.99%) sourced from Kurt J. Lesker. The ingots were flipped at least 5 times and remelted to ensure homogeneity. Ingots were encapsulated in quartz ampules purged with Ar gas and annealed at 800 °C for 24 h and then quenched in water at room temperature to preserve the desired microstructure. The ingots were sliced into 5 mm thick cylinders and mirror polished up to 1200 grit SiC paper finish. X-ray diffraction pattern analysis using a Bruker D2 Phaser diffractometer, which has a Cu anode ($k_α$, λ = 0.154 nm) using an accelerating voltage of 30 keV and a current of 10 mA. Chemical homogeneity was confirmed via energy dispersive spectroscopy (EDS) mapping on a FEI Quanta 650 field emission scanning electron microscope (SEM) with Oxford Instruments Aztec X-ManN EDS software.

Thermodynamic calculations based on the CALculation of PHAse Diagrams (CALPHAD) were performed using proprietary Thermo-Calc software using the TC-HEA3 database. Single-point equilibrium calculations were performed for the alloys at their relevant homogenization temperature utilizing their nominal compositions. Relevant properties (e.g., elemental activities) were recorded and utilized in ascertaining possible TEE mechanisms for Al. Calculations were performed at the homogenization temperature to reflect properties relevant to the equilibrium microstructure at that temperature, which was “frozen in” upon quenching. Equilibrium thermodynamic calculations performed at ambient temperatures

would be less accurate than the observed microstructure, as there is insufficient thermal energy at room temperature to allow lattice relaxation to the true equilibrium, properties of which CALPHAD would return. Thus, the appropriate procedure is to consider the microstructure formed at the homogenization temperature prior to quenching. Relevant properties (e.g., elemental activities) were recorded and utilized in ascertaining possible TEE mechanisms for Al.

Electrochemical testing was performed using either Gamry 600+ or Biologic SP-200 potentiostat and a three-electrode cell with a mercury-mercury sulfate reference electrode (E = 0.64 V vs. saturated hydrogen electrode (SHE)), a Pt mesh counter-electrode, and the alloy sample as the working electrode. All potentials in this work are reported vs. SHE. The exposed area was limited to 0.066 $cm^2$ using a Viton O-ring. Test solutions of 0.1 M $Na_2SO_4$ were prepared using sodium sulfate (Fischer Scientific) and titrated to a pH of 1 using sulfuric acid (stock Sigma Aldrich, 96%). A 3-step cathodic reduction was performed before each electrochemical test. These cathodic potential holds consisted of -0.76 V for 5 min, -1.26 V for 3 s, and -0.76 V for 1 min. Hydrogen gas bubbles formed on the sample surface during these steps were completely removed via frequent flushing of the solution using a pipette.

Firstly, anodic linear sweep voltammetry (LSV) experiments were performed after the cathodic reduction steps to determine the log(i)-E behavior of each alloy sample. The behavior of each sample was assessed from a scan range of -0.5 V to 1.2 V at a scan rate of approximately 0.52 $mV \cdot s^{-1}$. Next, independently, chronoamperometry (CA) experiments were conducted at 0.15 V for 300 s after performing the cathodic reduction steps in deaerated 0.1 M $H_2SO_4$. The number of dissolved alloy monolayers required to achieve passivity (h) was determined using the following relation:²⁴

$$h = \int_{0}^{t} i dt \cdot \left[ \frac{2 \cdot q \cdot (2 \cdot X_{Fe} + 3 \cdot X_{Cr} + 3 \cdot X_{Al})}{\sqrt{2} \cdot a^2} \right]^{-1} \quad [1]$$

where, i is the current density, q is the electronic charge, $X_M$ is the mole fraction of metal M in the alloy, and a is the lattice parameter. Valency of Fe, Cr, and Al was assumed to be 2+, 3+, and 3+, respectively. For simplicity, it was assumed that (110) planes of the BCC lattice are exposed.

Finally, potentiostatic film growth experiments were performed for 10 ks at two potentials: 0.15 V and 0.6 V, separately, and directly after performing the cathodic reduction steps. The two potentials were chosen to explore both the primary passive region at 0.15 V and the fully passive region at 0.6 V. Current density (i) along with the imaginary component of impedance ($-Z_{imag}$) was tracked as a function of time using an AC sinusoidal signal with an amplitude of 20 mV and a fixed frequency of 5 Hz. At this frequency, $-Z_{imag}$ is dominated by passive film oxide capacitance, enabling an estimation of its relative thickness, assuming a fixed dielectric constant.³⁸,³⁹ This experiment is hereby termed as single frequency electrochemical impedance spectroscopy (SF-EIS). At the end of every 10 ks hold, a broad spectrum EIS was conducted at the same DC potential using the same AC sinusoidal signal across a frequency range of 100 kHz to 5 mHz, measured every 8 points per decade. The EIS data were fit to an equivalent circuit model (ECM)⁴⁰ to represent the metal/oxide passive film/electrolyte interfaces. Furthermore, it should be noted that the low frequency limit of 5 mHz was selected to give information on corrosion protection, while the 5 Hz data was used to track relative oxide thickness.

Specific elemental dissolution current densities ($i_M$) of Fe, Cr, and Al were tracked of $FeCr_{10}$ and $FeCr_{10}Al_{16}$ alloys during CA experiments at 0.15 V for 2500 s (hereby termed as AESEC-CA) in deaerated 0.1 M $H_2SO_4$ using an AESEC apparatus described elsewhere.⁴¹ A cathodic reduction step at -0.76 V for 5 min was used to reduce the surface native oxides. The spectral signal of the dissolved species in the electrolyte was analyzed using an inductively coupled plasma atomic emission spectrometer (ICP-AES,

<table><caption>**Table I.** Alloy names and compositions of arc melted alloys after 24 h homogenization treatment at 800 °C determined by SEM-EDS in relation to its nominal composition in subscripts.</caption><thead><tr><th>Alloy name</th><th>Fe (at.%)</th><th>Cr (at.%)</th><th>Al (at.%)</th></tr></thead><tbody><tr><td>FeCr<sub>10</sub>Al<sub>4</sub></td><td>85.3</td><td>10.6</td><td>4.2</td></tr><tr><td>FeCr<sub>10</sub>Al<sub>6</sub></td><td>83.8</td><td>10.5</td><td>5.7</td></tr><tr><td>FeCr<sub>10</sub>Al<sub>8</sub></td><td>81.3</td><td>10.4</td><td>8.2</td></tr><tr><td>FeCr<sub>10</sub>Al<sub>10</sub></td><td>77.4</td><td>10.2</td><td>10.3</td></tr><tr><td>FeCr<sub>10</sub>Al<sub>16</sub></td><td>71.6</td><td>11.6</td><td>16.7</td></tr><tr><td>FeCr<sub>10</sub></td><td>89.3</td><td>10.8</td><td>—</td></tr><tr><td>FeCr<sub>16</sub></td><td>83.6</td><td>16.4</td><td>—</td></tr><tr><td>FeAl<sub>8</sub></td><td>91.06</td><td>—</td><td>8.94</td></tr><tr><td>FeAl<sub>16</sub></td><td>84.4</td><td>—</td><td>15.6</td></tr></tbody></table>
Horiba France), consisting of a polychromator with a focal length of 0.5 m and a monochromator with a focal length of 1.0 m. The Al signal at the wavelength of 167.02 nm was obtained using the monochromator to improve its detection limit. The emission intensities ($I_\lambda$) at characteristic wavelengths ($\lambda$) measured by the spectrometer for each element M was converted to concentration ($C_M$) in $nmol \cdot cm^{-3}$ using ICP-AES calibration method of using blank electrolyte, and standard electrolytes containing 0.2 ppm and 0.4 ppm of the elements to obtain a calibration curve.

$$C_M = (I_\lambda - I_\lambda^*)/k_{\lambda,M} \quad [2]$$

where $I_\lambda^*$ and $k_{\lambda,M}$ are the background intensity and the sensitivity factor for element M, respectively, for a given wavelength. The concentration was in turn, converted into current density ($i_M$):

$$i_M = n_M \cdot F \cdot f \cdot C_M/A \quad [3]$$

where, f is the electrolyte flow rate of $0.02 \ cm^3 \cdot s^{-1}$, A is the exposed flow cell surface area of $1 \ cm^2$, $n_M$ is the oxidation state of metal M as described previously for calculating h-values, and F is Faraday's constant. Total anodic current from dissolution ($i_{sum}$) was calculated as $\Sigma i_M$. Convoluted electrical current densities ($i_e^*$) were obtained from actual electrical current density ($i_c$), considering the electrolyte flow hydrodynamics between the electrochemical cell, and ICP-AES is elaborated elsewhere.41 Assuming Fe undergoes congruent dissolution in these alloys, hypothetical congruent dissolution current density ($i_M^{cg}$) for Cr and Al were obtained using the relation:

$$i_M^{cg} = \frac{n_M \cdot X_M}{n_{Fe} \cdot X_{Fe}} \cdot i_{Fe} \quad [4]$$

where, $X_M$ is the mole fraction of element M (Cr or Al). The difference between the congruent and actual dissolution rate indicates the element's enrichment rate. The differences between the area under the curves $i_M^{cg}$ and $i_M$ for a given interval of potential or time would indicate the elemental enrichment $\Theta_M$ in $C \cdot cm^{-2}$ at the surface. The percent elemental enrichment ($\Theta_M'$) is normalized to the total congruent dissolution of the element as described below:

$$\% \Theta_M' = \left\{ 1 - \frac{\int_a^b i_M dt}{\int_a^b i_M^{cg} dt} \right\} \times 100\% \quad [5]$$

where, $a = 30$ s and $b = 2500$ s are the lower and upper limits of time intervals. For an element M (Cr or Al), $\Theta_M' > 0$ indicates surface enrichment, while $\Theta_M' < 0$ signifies surface depletion.

Chemical compositions of the oxide passive films formed after the 10 ks potentiostatic hold experiments at 0.15 V and 0.6 V in 0.1 M $H_2SO_4$ were analyzed via X-ray photoelectron spectroscopy (XPS). Samples were rinsed in DI water afterward to eliminate any sulfate salt residue on the surface and carried in a $N_2$ flushed bag for less than 5 min before introducing them into the XPS chamber. A PHI Versa Probe-III XPS analyzer, calibrated to Au 4f$_{7/2}$ at 84.00 eV, was used to detect the emitted photoelectrons. Core shell spectra of O 1s, Fe 2p$_{3/2}$, Cr 2p$_{3/2}$, and Al 2p were collected using an Al X-ray beam with a spot dimension of 100 x 100 µm, power 100 W, pass energy 26 eV, and 0.05 eV step size. Carbon 1s at 284.8 eV was used as the standard to correct for peak shifts due to surface charging. Spectral deconvolution peak fitting of the core-shell spectra was performed following the Biesinger et al.42 for Fe 2p$_{3/2}$, and Cr 2p$_{3/2}$, with an acceptable deviation (< 0.75 eV) of peak positions, see Tables S4 and S5 in the supplementary material. Details of the spectral deconvolution peak fitting to obtain cationic fractions are described elsewhere.29,43

## Results

**Alloy chemical composition and phase analysis.**—Table I shows the targeted and measured compositions of the alloys studied in this

<figure><img src="image_6.png" /><figcaption>**Figure 1.** XRD pattern analysis of FeCr$_{10}$Al$_x$, FeAl$_x$, and FeCr$_x$ alloys studied in this work after homogenization heat-treatment at 800 °C for 24 h.</figcaption></figure>

work obtained from SEM-EDS analyses. All compositions are given in atomic percent and disregard any contaminants such as silicon or carbon. These compositions were also further validated by X-ray fluorescence and ICP-OES analysis of digested alloys in Aqua Regia (3 M HCl + 1 M $HNO_3$). These results are reported in Tables S1-2 of the supplementary material. Hereafter, the alloys will be referred to following their names listed in Table I. XRD pattern analysis for all the alloys is shown in Fig. 1, such that all alloys showed only the peaks representing BCC crystal structure following the signature peaks similar to that of BCC Fe,44 indicating all are single phase after the solutionization heat-treatment.

**Anodic passivity from log(i) vs. E plots.**—Figure 2a shows the log(i) vs. E curves obtained from LSV experiments performed in deaerated acidified 0.1 M $Na_2SO_4$. An active-to-passive behavior of all the alloys was observed. Active corrosion occurs from the alloy corrosion potential around -0.4 V until a current density peak to attain passivity, defined as critical current density $i_{crit}$, is reached at different potentials ($E_{crit}$) for the alloys. All alloys tested showed similar active dissolution rates before achieving $i_{crit}$. Passive current density ($i_{pass}$) is defined as the current density of the passivated alloy surface. FeAl$_x$ alloys showed a trend of decreasing values of $i_{crit}$, $i_{pass}$ as well as $E_{crit}$ with increasing Al content from 4, 16, and 24 at.%. However, both their $i_{crit}$ and $i_{pass}$ values are more than 5 times, and compared to those of FeCr$_{10}$ alloy, suggesting even 24 at. % Al is not enough to compete with 10 at. % Cr in Fe. Within the FeCr$_{10}$Al$_x$ series, a clear decrease in $i_{crit}$ as well as $E_{crit}$ values was observed, ranging from the FeCr$_{10}$ binary alloy with sequentially increasing Al content from 4 to 16 at. %, approximately matching the behavior of FeCr$_{16}$. This indicates the beneficial effect of alloying Al, significantly improving the ability of FeCr$_{10}$ alloy to passivate.

To compare results and place in the context of the literature, $i_{crit}$ is plotted as a function of Al content in the alloy in Fig. 2b for binary FeAl$_x$ alloys compared to FeCr$_{10}$Al$_x$. It was observed that the $i_{crit}$ values for all ternary alloys are significantly lower than their binary FeCr$_x$ and FeAl$_x$ counterparts. It is also interesting that they seem to approach a critical threshold of an $i_{crit}$ value of 2.5 $mA \cdot cm^{-2}$, comparable to that of the FeCr$_{17}$ and FeAl$_{30}$ alloys. This indicates Al addition to FeCr$_{10}$ alloy may not be equivalent to “Cr” addition and may require around 16 at. % Al to achieve the same passive formation behavior when more than 7 at. % Cr is added.

**h-values and thermodynamic activity of Cr upon Al addition.**—

Figure 3 illustrates the decrease in h-values in monolayers required by FeCr$_{10}$Al$_x$ alloys with increasing Al content. Here, h determines the number of monolayers dissolved during active-passive transition
<figure><img src="image_7.png" /><figcaption>**Figure 2.** (a) Log(i) vs. E curves of FeAlₓ, FeCrₓ as well as FeCr₁₀Alₓ alloys obtained from anodic LSV experiments in deaerated 0.1 M Na₂SO₄ titrated to pH 1. (b) Comparison of *i*<sub>crit</sub> values plotted vs. at.% Al obtained from LSV of all the FeCr₁₀Alₓ tested in this work with the same obtained from literature for FeAlₓ<sup>26</sup> alloys in 0.0126 M H₂SO₄ and FeCrₓ<sup>24</sup> alloys in 0.1 M H₂SO₄.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>**Figure 3.** Number of dissolved monolayers to achieve passivity (h) obtained from CA experiments at 0.15 V in 0.1 M H₂SO₄ plotted vs. Al content in FeCr₁₀Alₓ alloys. The gray solid line denotes the h-value of 15 monolayers of FeCr₁₆ in deaerated 0.1 M H₂SO₄.<sup>24</sup> Red and blue colors in the color bar indicate large and small values of Cr activity in the BCC phase FeCr₁₀Alₓ at 800 °C, obtained from ThermoCalc.</figcaption></figure>

to uncover enough of the passivating elements to form a passivating film according to percolation.²⁴ The potential of 0.15 V was chosen for the CA experiments as it was the lowest anodic potential above all alloys' active-to-passive transition, see Fig. 2a. The FeCr₁₀ alloy showed a significantly higher *h*-value of 950 monolayers. Without Cr, FeAlₓ does not passivate at 0.15 V and can be expected to require more than 950 monolayers.⁴⁵ With increasing Al additions to FeCr₁₀, there is a clear trend of decreasing the number of monolayers required to passivate to as low as 15 monolayers such that it is comparable to those required by FeCr₁₆ alloy.²⁴ Further, just a small addition of 4 at.% Al resulted in a ten-fold lower *h*-value, indicating

that even a small presence of Al is beneficial in supporting the Cr-lean alloy to passivate.

To provide insight into Al's role in promoting Cr-based passive film formation, the data in Fig. 3 are colored by the relative activity of Cr in solid solution calculated using the CALPHAD methodology, with red indicating higher activity and blue indicating lower. As predicted by CALPHAD, increasing Al content at fixed Cr amounts (with Fe as balance) led to relatively higher Cr activity in the alloy. Activity is commonly interpreted as an "effective concentration," so Al additions are suggested to make the alloy behave as though it contains more Cr than is nominally present in the bulk. This trend, in which increasing Al simultaneously increases Cr activity and decreases *h*, suggests that Cr activity may be a strong indicator of overall passivation charge efficiency (characterized by *h*). Moreover, this connection between a readily computable thermodynamic property and a measured indicator of passivation could provide a path toward high-throughput design of stainless alloys, using CALPHAD-informed models to optimize alloy composition for enhanced passivator performance.³²

**Potentiostatic passive film growth via SF-EIS.**—Figures 4a–4b illustrate the *i* and −*Z*<sub>imag</sub> data obtained from SF-EIS experiments at two applied potentials, 0.15 V and 0.6 V, respectively. This test tracked the electric current density associated with the oxidation of all elements and the relative film thickness given by −*Z*<sub>imag</sub>. In general, at 0.15 V, all alloys demonstrated a slower kinetics for passivity compared to the 0.6 V, due to their smaller anodic overpotential (driving force). At 0.15 V, additions of Al ranging from 4–16 at.% to a FeCr₁₀ alloy promoted forming thicker films, as the FeCr₁₀ indicated any film formed showed a negligible rise in −*Z*<sub>imag</sub> value with time. The FeCr₁₀Al₆ alloy forms a film of similar relative thickness (i.e., similar −*Z*<sub>imag</sub> value) to FeCr₁₆ but reaches a low constant current density while FeCr₁₆ continues to decrease its current density through the end of the test. The current of FeCr₁₀Al₁₆ alloy rapidly decays, and this alloy forms a passive film which has the greatest change in relative thickness from 0–100 s. It exhibits a −*Z*<sub>imag</sub> value larger than FeCr₁₆. At the conclusion of the test for
<figure><img src="image_9.png" /><figcaption>**Figure 4.** Current density (i) and imaginary impedance (−Z<sub>imag</sub>) vs. time data of FeCr<sub>10</sub>Al<sub>x</sub>, select FeAl<sub>x</sub> and FeCr<sub>x</sub> alloys obtained from SF-EIS experiments performed for 10 ks at (a) 0.15 V and (b) 0.6 V. Electrolyte used was deaerated 0.1 M Na<sub>2</sub>SO<sub>4</sub> titrated to pH 1. Note: FeAl<sub>8</sub> was not tested because of its complete lack of passive film and extreme noisy response.</figcaption></figure>

10 ks, the current density dropped as Al increased from 4 to 16 at.%.
Note that FeCr<sub>10</sub>Al<sub>16</sub> has a much greater –Z<sub>imag</sub> than the other
compositions, yet has a similar passive response seen in FeCr<sub>16</sub>.

**Passive film impedance from broad spectrum EIS.**—Figure 5a illustrates the Bode EIS plots obtained at 0.15 V after SF-EIS experiments for 10 ks for select alloys. A higher impedance modulus (|Z|) value at 5 mHz frequency is indicative of a more protective passive film. At 0.15 V, FeAlₓ alloys were omitted due to their lack of any passive formation tendency at this potential, refer to logi-E behavior shown in Fig. 2a. FeCr₁₀ alloy showed the smallest value of |Z| at 5 mHz of approximately 0.005 kΩ·cm². Systematic addition of 4 to 16 at.% Al in FeCr₁₀ increased this value by a few orders of magnitude, such that FeCr₁₀Al₄ exhibited a value of 1.5 kΩ·cm², FeCr₁₀Al₁₀ a value of 20 kΩ·cm², and FeCr₁₀Al₁₆ a value of 200 kΩ·cm². This indicates the significant benefit of Al alloying on the passive film impedance magnitude in the primary passive region. Figure 5b illustrates the Bode EIS plots after SF-EIS experiments at 0.6 V for selected alloys. All Cr containing FeCr₁₀Alₓ as well as FeCrₓ alloys demonstrated a higher |Z| value of 100 kΩ·cm² at 5 mHz, compared to 10 kΩ·cm² of FeAl₁₆ alloy. However, within FeCr₁₀Alₓ and FeCrₓ alloys, insignificant differences in |Z| at 5 mHz were observed, suggesting neither benefit nor degradation in steady-state passivity with Al addition at 0.6 V.

The EIS data for some selected alloys were fit using an ECM rigorously reported for fitting passive film impedance data.⁴⁰ Here, R_s, R_{fle}, and R_f represent the resistance from the solution, film/ electrolyte interface, and the passive film, respectively. Q represents a constant phase element (CPE), while W the Warburg diffusional impedance for ionic transport kinetics. Polarization resistance (R_p) of the passive films is herein defined as R_{fle} + R_p, in kΩ·cm². Individual ECM element parameters along with R_p for FeCr₁₆, FeCr₁₀Al₁₀, and FeCr₁₀Al₁₆ are listed in Table II. R_p followed the same trend for |Z|² at 5 mHz with FeCr₁₀Al₁₆ demonstrating a value of 269 kΩ·cm² compared to 13.4 kΩ·cm² for FeCr₁₆. The ECM circuit parameters post-fitting the EIS data obtained at 0.6 V demonstrated a large but similar R_p value for both FeCr₁₀Alₓ and FeCrₓ alloys, as shown in Table S3 of the supplementary material.

**Tracking elemental dissolution in-operando via AESEC.**—
Elemental dissolution currents of Fe (Fe → Fe²⁺), Cr (Cr → Cr³⁺), and Al (Al → Al³⁺) during the anodic AESEC-CA experiments of FeCr₁₀ and FeCr₁₀Al₁₆ performed at 0.15 V are shown in Fig. 6. The convoluted current density, i* is also shown to account for the broadening of the spectroscopic signal caused by the hydrodynamics of the flow cell.⁴¹ It allows direct comparison between the elemental dissolution current signals and the convoluted electrical signal. Alloy FeCr₁₀Al₁₆ showed approximately 10 times lower iₑ compared to FeCr₁₀ throughout the experiment. For FeCr₁₀, the total elemental dissolution current density (iₐsum) representing the sum of all Fe and Cr dissolution currents was slightly lower than iₑ, indicating a near Faradaic dissolution process. In contrast, the dissolution current was significantly below the congruent expectation for FeCr₁₀Al₁₆, indicating passive film formation, i.e., unaccounted charge due to the passive film formation.⁴¹ Compared to FeCr₁₀Al₁₆, an approximate 80 times greater equivalent Fe and Cr dissolution current density was observed for FeCr₁₀, which increased with time. In contrast, for the case of FeCr₁₀Al₁₆, Fe and Cr dissolution showed a peak value of 15 µA·cm⁻² and 6 µA·cm⁻², respectively, and then decreased with time. A similar behavior was observed for Al in FeCr₁₀Al₁₆, which showed a peak value of 2 µA·cm⁻². Assuming congruent dissolution of Fe, between 20–2500 s, both alloys demonstrated surface enrichment (Θ'ₘ) of Cr since i<sub>Cr</sub> < i<sub>Cr</sub><sup>cg</sup> such that it was only 1.1% for FeCr₁₀ but 22.7% for FeCr₁₀Al₁₆. Similarly, Al showed a surface enrichment of 46.3% on the FeCr₁₀Al₁₆ alloy. This suggests that alloying with Al leads to increased enrichment of Cr³⁺ as well as Al³⁺ cations on the surface as the surface re-passivates.

**Oxide passive film chemistry via XPS.**—Cationic species such as metallic oxides and hydroxides present in the potentiostatically grown passive films of all the FeCr<sub>10</sub>Al<sub>x</sub> alloys were analyzed using XPS. Figure 7 demonstrates a representative deconvolution peak fitting of the O 1s, Al 2p, Cr 2p<sub>3/2</sub>, Fe 2p<sub>3/2</sub> core shell spectra. The O 1s spectra indicated the presence of metallic oxides and hydroxides present in the passive film, as expected in an aqueous electrolyte.⁴³ A majority of Fe(II) and some Fe(III) oxides were observed with a
<figure><img src="image_10.png" /><figcaption>**Figure 5.** Bode impedance (|Z|) and phase (Z<sub>phz</sub>) plots obtained from EIS experiments on select FeCr<sub>10</sub>Al<sub>x</sub>, FeCr<sub>x</sub>, and FeAl<sub>x</sub> alloys at applied potential of (a) 0.15 V and (b) 0.6 V, both after 10 ks of SF-EIS experiments at the same potential. The data is shown as open circles. (c) Equivalent circuit model used to fit selected EIS data, shown as solid lines in sub Figs. (a) and (b). The electrolyte used was deaerated 0.1 M Na<sub>2</sub>SO<sub>4</sub> titrated to pH 1. Note: Data for FeAl<sub>x</sub> alloys are not shown at 0.15 V due lack of surface passivity and thus, a clean EIS response. The same is the case for FeAl<sub>8</sub> at 0.6 V.</figcaption></figure>

<table><caption>**Table II.** Values of ECM circuit elements after fitting EIS data obtained for selected alloys at 0.15 V after 10 ks in deaerated 0.1 M Na<sub>2</sub>SO<sub>4</sub> titrated to pH 1.</caption><thead><tr><th>Alloy</th><th>R<sub>s</sub><br>(Ω·cm<sup>2</sup>)</th><th>R<sub>fe</sub><br>(Ω·cm<sup>2</sup>)</th><th>Q<sub>fe</sub><br>(μS·cm<sup>-2</sup>·s<sup>-1</sup>)</th><th>a<sub>fe</sub></th><th>R<sub>f</sub><br>(kΩ·cm<sup>2</sup>)</th><th>W<br>(Ω·cm<sup>2</sup>·s<sup>-1/2</sup>)</th><th>Q<sub>f</sub><br>(μS·cm<sup>-2</sup>·s<sup>-1</sup>)</th><th>a<sub>f</sub></th><th>R<sub>p</sub><br>(kΩ·cm<sup>2</sup>)</th></tr></thead><tbody><tr><td>FeCr<sub>16</sub></td><td>6.2</td><td>1.7</td><td>3.93</td><td>0.92</td><td>1.6</td><td>6667.9</td><td>129.8</td><td>0.87</td><td>181.4</td></tr><tr><td>FeCr<sub>10</sub>Al<sub>10</sub></td><td>3.9</td><td>7.6</td><td>0.3</td><td>0.92</td><td>13.3</td><td>957.2</td><td>107.3</td><td>0.88</td><td>13.4</td></tr><tr><td>FeCr<sub>10</sub>Al<sub>16</sub></td><td>3.6</td><td>4.5</td><td>0.3</td><td>0.93</td><td>269.1</td><td>11640.4</td><td>56.1</td><td>0.89</td><td>269.1</td></tr></tbody></table>

strong Fe metal peak, probed from the alloy beneath the passive film, which is a few nanometers thick. The cation fractions obtained from peak fitting are shown in Fig. 8. An element is enriched when it is present in a higher fraction of its cations in the passive film than the fraction of the element in the alloy; otherwise, it is depleted.⁴⁶ At both potentials, Fe(II, III) cations are the dominant species. At 0.6 V, a consistent decrease in the Fe(II, III) cation fraction is observed with increasing Al content from 4 to 16 at.% in FeCr<sub>10</sub> alloy. A strong enrichment of Al(III) is observed, while Cr(III) is slightly enriched for all the FeCr<sub>10</sub>Al<sub>x</sub> alloys. At 0.15 V, both Cr(III) and Al(III) cations are enriched with increasing Al addition, while Fe(II, III) cations are depleted in the passive film. This is consistent with the AESEC-CA results of positive values of %Θ<sub>Cr,Al</sub> at for FeCr<sub>10</sub>Al<sub>16</sub> alloy at 0.15 V, as well as with the displayed large gains in corrosion protection property of passive film impedance at 0.15 V, both

described above. Overall, the cation fractions suggest a TEE on Al addition is present with increasing passivating cation fractions.²⁹,³⁰

## Discussion

Two binary alloys, FeCr<sub>10</sub> and FeCr<sub>16</sub>, functioned as baselines for the ternary alloy series. FeCr<sub>16</sub> contains enough Cr to form a continuous protective film in acidified sulfate solution, as seen by the large passive region in Fig. 2.²³ FeCr<sub>10</sub>, on the other hand, had a very minimal passive region. Recall that 10 at.% Cr is below the critical Cr often cited with respect to Fe-Cr alloys but also seen in other alloys such as Ni-Cr and Co-Cr.⁴⁷ The main findings of this work are: (i) Addition of Al enhanced the aqueous passivation behavior of FeCr<sub>10</sub> alloys in forming a continuous protective passive film by introducing a wider passivity window, lower h-values, as
<figure><img src="image_11.png" /><figcaption>**Figure 6.** Equivalent dissolution current densities ($i_M$) of Fe, Cr, Al, convoluted ($i_e^*$), and actual ($i_e$) electrical current density obtained from AESEC-CA experiments of FeCr$_{10}$ and FeCr$_{10}$Al$_{16}$ alloys at 0.15 V for 2500 s in deaerated 0.1 M $H_2SO_4$. Note the difference in the y-axis scale from left to right in the subfigures. Here, $\Theta'_{Cr}$ and $\Theta'_{Al}$ represent the percentage of Cr and Al enriched on the surface between 20 - 2500 s.</figcaption></figure>

well as higher $|Z|$ value at 5 mHz after 10 ks of passive film growth at 0.15 V, and (ii) Al additions introduced TEE type behavior analogous to that at elevated temperatures (>1000°C). However, here, both Cr(III) and Al(III) cationic enrichment was observed, via both XPS based cation fractions as well as *in-operando* tracking of dissolution rates via AESEC, both at 0.15 V.

A major positive shift in passivity was brought with just 4 at.% of Al addition to FeCr$_{10}$. FeCr$_{10}$Al$_x$ alloys exhibited a nearly identical E-log (i) behavior at more anodic potentials above 0.3 V vs. SHE similar to that of FeCr$_{16}$ (see Fig. 2a). FeCr$_{10}$Al$_4$ showed marked improvement from FeCr$_{10}$ in both long-term passive film impedance as well at 0.15 V (see Figs. 4a and 5a). The oxide film thickness increases, as well as produced a drop in current density that levels out about 100 s in each experiment. While it is not close to the performance of FeCr$_{16}$, it is noticeably much better than the performance of FeCr$_{10}$, which did not grow a noticeably thick film and only dropped about one order of magnitude in current density during single-step SF-EIS experiments. Similar behavior to this observation in the case of 4 at.% Al, the same trend was seen for 6, 8, and 16 at.% Al added FeCr$_{10}$ alloys. Hio et al. reported a comparable study on the passivity of FeCrAl alloys, except that their work only involved the addition of about 6 at.% Al and did not vary the Al content.<sup>48,49</sup> When evaluated over a range of Cr contents at one Al concentration in $H_2SO_4$ solution, a large effect was observed upon Al additions when Cr was fixed 10 at.% Cr.<sup>48,49</sup> Further, in the work of Hio, there was no single step passivation impedance spectroscopy or XPS analysis. In this study, the same behavior can be seen to extend to long term passive film impedance and oxide thickness, as well as by decreasing *h*-values.

One of the most notable results from the SF-EIS is that Al plays a more significant role in the primary passive region at 0.15 V vs SHE than in the fully passive region at 0.6 V vs SHE. In the fully passive region, there is noticeable overlap in passive current density between all of the ternary alloys, whereas the primary passive region shows the impact of small variations in Al. This could be in part due to Al's role of suppressing Fe oxide formation more than promoting $Cr_2O_3$ formation, as observed from XPS cation fractions (see Fig. 8). In the early-stage passivation, Al may suppress Fe dissolution along with oxide formation in order to better facilitate the growth of both Cr and Al oxides. This is strongly observed from SF-EIS experiments at 0.15 V. Here, dissolution of Fe along with Cr and Al was significantly reduced, while Cr(III) and Al(III) cations were enriched on the surface (see Figs. 6 and 8). Effectively, Al may pull the oxygen into its $Al_2O_3$ structure rather than forming less protective Fe oxides or mixed oxides, functioning as an “O getter”.<sup>6,50</sup> Since $Cr_2O_3$ and $Al_2O_3$ have the same crystal structure, and valency, it is possible to assume that this also facilitates $Cr_2O_3$ growth utilizing a a scaffold of $Al_2O_3$ that can form first.

Al has been shown to increase the thermodynamic activity of Cr, which may indicate that there is effectively more Cr available to form its $Cr_2O_3$ rich passive film due to the presence of Al. Such alterations have been reported in other systems, an example being where Al activity was used as a metric for achieving Al$_2O_3$-rich scales on Nb-Si-Ti-Al-Hf alloys at elevated temperatures.<sup>51</sup> Furthermore, from first principles calculations, Al in BCC Fe-Cr-Al has been predicted to introduce clustering type chemical short-range order (CSRO).<sup>30</sup> This is due to its strong tendency for ordering with Fe (more Fe-Al bonds),
<figure><img src="image_12.png" /><figcaption>**Figure 7.** XPS deconvolution peak fits of the O 1s, Fe 2p<sub>3/2</sub>, Cr 2p<sub>3/2</sub>, and Al 2p core-shell spectra of FeCr<sub>10</sub>Al<sub>6</sub> alloy passive film obtained after 10 ks of SF-EIS experiment at 0.15 V in deaerated 0.1 M H<sub>2</sub>SO<sub>4</sub>. Metallic peaks are shown in pink, oxides are shown in cyan, and hydroxides are shown in olive green. The Cr 3s peak in the Al 2p region is shown in orange.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Figure 8.** Passive film cation fractions for Fe, Cr, and Al determined from deconvolution peak fitting of Fe 2p<sub>3/2</sub>, Cr 2p<sub>3/2</sub>, and Al 2p XPS core-shell spectra of FeCr<sub>10</sub>Al<sub>x</sub> alloys. The XPS spectra were obtained after SF-EIS experiments performed for 10 ks at 0.6 V and 0.15 V in deaerated 0.1 M H<sub>2</sub>SO<sub>4</sub>. Here, circles and solid lines represent data from 0.6 V, and triangles and dashed lines from 0.15 V.</figcaption></figure>

which increases the probability of finding more Cr-Cr clustered pairs in the first few nearest neighbor shells.⁴⁵ This can translate to more Cr-Cr pairs available to form Cr<sub>2</sub>O<sub>3</sub> during the primary passivation stages of alloy passive film formation, as was observed from AESEC-CA and XPS analyses at 0.15 V (see Figs. 6 and 8). This increase in more available Cr atoms in the alloy can be another way in which Al facilitates enhanced passivity.

One of the classical TEE mechanisms posits that the passive film develops a complex structure, perhaps with Al(III) oxide or hydroxide forming an early layer, which limits Fe(II) and Fe(III) transport and serves as scaffolding for Cr(III). Since Al<sub>2</sub>O<sub>3</sub> and Cr<sub>2</sub>O<sub>3</sub> have the same crystal structure and similar characteristics, it may be possible that a complex oxide or solute trapped oxide that is not a spinel is responsible for improved passivation in FeCrAl alloys. It cannot be ruled out that solute trapping plays a role, nor can spinels be entirely dismissed. Future work involves 3-dimensional atom probe tomography to further the understanding of the structure of these passive films.⁵²

## Conclusions

In this work, low Cr containing FeCrAl alloys were studied to explore the *third-element effect* of Al addition on the aqueous passivity of these alloys. A suite of electrochemical parameters was utilized, such as *i*<sub>crit</sub>, *i*<sub>pass</sub>, *h*-values, and |*Z*| at 5 mHz, surface-sensitive XPS-based elemental cation fractions, and *in-operando* spectroelectrochemical parameters including *i*<sub>M</sub> and %Θ'<sub>Cr,Al</sub>. The
following conclusions were observed in this study which was
conducted using sulfuric acid:

* Additions as low as 4 at.% of Al to 10 at.% Cr containing Fe-base alloys, significantly improved corrosion resistance characterized by lower $E_{crit}$, $i_{crit}$ and $h$-values in acidified sulfate solutions. Larger additions up to 16 at.% Al to an $FeCr_{10}$ alloy matched the passive behavior of $FeCr_{16}$.

* Al-containing alloys also resulted in thicker (greater $-Z_{imag}$) and more protective passive films developing over time, with greater $|Z|$ value at 5 mHz and polarization resistance in acidified sulfate solutions. The TEE of Al was observed in achieving passivity of a Cr lean $FeCr_{10}$ alloy by enriching both the presence of protective Cr(III) and Al(III) cations in the passive films, while also significantly reducing the dissolution kinetics of Fe, along with Cr and Al in the electrolyte.

* The impact of Al with respect to the TEE has been explored in this article. There is clear evidence of the promotion of passive film growth in $FeCrAl$ alloys with increasing Al content, even when Cr is below the classical threshold of 12 at.%. Al(III) oxides are present and likely play an active role and can be observed in the long-term passive oxide, as explored by XPS and AESEC.

Acknowledgments

The authors gratefully acknowledge partial funding from the Office of Naval Research (ONR) through the Multidisciplinary University Research Initiative (MURI) program (award: N00014-20-1-2368) with program manager Dr. D. Shifler. Thanks to Kaitlyn L. Anderson for aiding with some of the electrochemical data acquisition. C.L. also acknowledges the U.S. Nuclear Regulatory Commission University Nuclear Leadership Program (NRC UNLP) for additional funding of this project (fellowship award: 31310023M0034). D.S. extends his gratitude to the UVA-SEAS Endowed Olsen Graduate Fellowship (2024) for partial funding of the AESEC work at CNRS-LISE, Sorbonne University, Paris. J.H. and AESEC experiments conducted by D.S. were also supported by the French government's "France 2030" initiative through the PEPR-DIADEM (Priority Research Programs and Equipment - Integrated Devices for Accelerating the Deployment of Emerging Materials) program, managed by the French National Research Agency (Agence Nationale de la Recherche, ANR), n°ANR-23-PEXD-0006. The authors also acknowledge the University of Virginia NMCF for utilizing their PHI Versaprobe III XPS, Bruker D2 XRD, Epsilon III XRF, and Quanta 650 SEM.

ORCID

Catherine S. Lynch https://orcid.org/0009-0007-7788-3373

Debashish Sur https://orcid.org/0000-0002-7954-9949

Peter F. Connors https://orcid.org/0009-0009-2341-0081

Benjamin W. Y. Redemann https://orcid.org/0000-0001-8863-4363

Junsoo Han https://orcid.org/0000-0001-9295-2954

Elizabeth J. Opila https://orcid.org/0000-0001-5540-7084

John R. Scully https://orcid.org/0000-0001-5353-766X

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