# Atomic structure, chemical composition, and oxidation behavior of the (001) surface of Ni 2 In 3 intermetallic catalyst

Ahowd Youssef Alfahad, Vipin Kumar Singh, Emilie Gaudry, Julian Ledieu, Vincent Fournée, Peter Gille, Adrian Gardner, Paul Chalker, Ronan Mcgrath, Hem Raj Sharma

► To cite this version:

Ahowd Youssef Alfahad, Vipin Kumar Singh, Emilie Gaudry, Julian Ledieu, Vincent Fournée, et al. Atomic structure, chemical composition, and oxidation behavior of the (001) surface of Ni 2 In 3 intermetallic catalyst. Physical Review Materials, 2025, 9 (3), pp.035801. 10.1103/PhysRevMateri-als.9.035801. hal-05046042

HAL Id: hal-05046042

https://hal.science/hal-05046042v1

Submitted on 10 Nov 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
# Atomic structure, chemical composition, and oxidation behavior of the (001) surface of Ni₂In₃ intermetallic catalyst

Ahowd Youssef Alfahad,¹,² Vipin Kumar Singh,³ Emilie Gaudry,³ Julian Ledieu,³ Vincent Fournée,³ Peter Gille,⁴ Adrian M. Gardner,⁵ Paul R. Chalker,⁵ Ronan McGrath,¹ and Hem Raj Sharma¹,*,

¹Surface Science Research Centre and Department of Physics, University of Liverpool, Liverpool L69 3BX, United Kingdom

²Department of Physics, College of Science and Humanities–Jubail, Imam Abdulrahman Bin Faisal University, Jubail 35811, Saudi Arabia

³Institut Jean Lamour, CNRS-Université de Lorraine, Campus Artem 2 allée André Guinier BP 50840 54011 Nancy Cedex, France

⁴Department of Earth and Environmental Sciences, Ludwig-Maximilians-Universität, München, Germany

⁵Low Energy Ion Scattering Facility, Geroge Holt Building, University of Liverpool, Liverpool L69 3GB, United Kingdom

<figure><img src="image_2.png" /></figure>

(Received 27 September 2024; revised 21 January 2025; accepted 5 March 2025; published 24 March 2025)

Intermetallic compounds have shown promising catalytic properties, exhibiting better selectivity and thermal stability than their metallic counterparts. Understanding their surface properties is crucial for optimizing catalytic performance. Here, we present an investigation of the atomic structure, chemical composition, and oxidation behavior of the previously unexplored binary intermetallic catalyst, Ni₂In₃, using various complementary techniques, including low-energy electron diffraction (LEED), scanning tunneling microscopy (STM), x-ray photoelectron spectroscopy (XPS), and low-energy ion scattering (LEIS), which provides elemental information from the topmost surface layer. We also calculated the surface energy of several possible bulk-terminated surfaces by density functional theory (DFT) to investigate surface stability. The Ni₂In₃(100) surface was prepared by ion bombardment and annealing under ultrahigh vacuum conditions. We find that ion bombardment preferentially removes indium (In), while annealing promotes In segregation to the surface. After annealing, the surface exhibits a (1 × 1) LEED pattern. The surface develops a three-domain c(2√3 × 4) rectangular structure upon exposure to ambient hydrogen present in the experimental chamber or hydrogen dosed from an external source, as evidenced by both STM and LEED. Both nickel (Ni) and In oxidize upon exposure to molecular oxygen. Based on STM and LEIS results along with surface energy calculations, we reveal that the surface planes contain In atoms bonded with Ni, the catalytically active constituent, suggesting that the surface may be stabilized by intermetallic bonding.

DOI: 10.1103/PhysRevMaterials.9.035801

## I. INTRODUCTION

Intermetallic compounds (IMCs) are composed of two or more elements with specific stoichiometry, with each constituent element occupying specific sites within the crystal structure. These compounds show potential applications in various areas, one of which is heterogeneous catalysis [1–4]. IMCs offer low-cost and environmentally friendly alternatives to metal catalysts. For example, Al₁₃Fe₄ IMC can replace Pd catalysts in semi-hydrogenation reactions [5]. Moreover, the electronic properties of IMCs can be tuned for catalytic purposes by altering chemical composition [6,7].

It is believed that the spatial separation of catalytically active sites in IMCs (referred to as site isolation) leads to higher selectivity and improved thermal stability compared to their metallic counterparts [8]. The concept of site isolation

has been experimentally tested on Pd-Ga intermetallics (Pd₃Ga₇, PdGa, and Pd₂Ga) for hydrogenation reactions [6,9–11]. In the Pd-Ga crystal structure, Pd is solely surrounded by Ga, isolating Pd atoms from each other. Additionally, the electronic structure of Pd-Ga around the Fermi level is strongly modified compared to elemental Pd. It was therefore concluded that electronic modification and isolation of active Pd sites result in the superior catalytic properties of Pd-Ga.

Similarly, Tsai et al. studied the relationship between catalytic activity and the valence electron density of states of isostructural ZnPd, ZnPt, and ZnNi [12]. They found that the valence electron density of states of ZnPd is similar to that of elemental Cu, unlike that of the other two compounds, and thus PdZn can potentially replace Cu catalysts in the steam reforming of methanol. It was suggested that the ZnPd bonds formed at surfaces may contribute to the excellent catalyst properties of ZnPd [13].

The In-Ni IMCs, the system of the current study, are found to be good environmental catalysts. These compounds can be used in the dry reforming of methane, a reaction that utilizes two greenhouse gases—carbon dioxide and methane—to produce other useful compounds [11]. These are also promising catalysts for the hydrogenation of carbonyl compounds

<footnote>*Contact author: H.R.Sharma@liverpool.ac.uk</footnote>

<footnote>Published by the American Physical Society under the terms of the Creative Commons Attribution 4.0 International license. Further distribution of this work must maintain attribution to the author(s) and the published article's title, journal citation, and DOI.</footnote>
[14]. Li *et al.* synthesized nanoparticles of various sizes and compositions of Ni-In IMCs (Ni₃In, Ni₂In, NiIn, and Ni₂In₃) catalysts using chemical methods and examined the hydrogenation reactions. They found that these compounds exhibit excellent catalytic activity and selectivity depending on the chemical composition and particle size. The NiIn and Ni₂In₃ nanoparticles perform better than the other species. Again, the superior performance was attributed to the isolation of the Ni sites and modification of electronic properties by electron transfer from In to Ni. Ni-based catalysts are also widely used for oxidation reactions, including the reduction of carbon monoxide [15], the conversion of methane to syngas [16–18], and the synthesis of industrial chemicals [19].

Real-world catalysts are very complex, making it challenging to gain insight into a fundamental understanding of catalytic behavior, and catalytic activity is usually optimized by trial and error methods. Therefore, simple and well-defined model systems are used to investigate the physical parameters influencing catalytic activity. Often, single crystal samples with surfaces cut along known crystallographic faces are used as model catalysts, and surface science techniques are employed under ultrahigh vacuum (UHV) conditions to characterize their atomic and electronic properties and their behavior towards the reactants of the respective chemical reactions [20].

In this paper, we present a surface study of a single crystal Ni₂In₃ intermetallic catalyst. The atomic structure, chemical composition, and oxygen adsorption behavior of the (001) surface using various techniques including low-energy electron diffraction (LEED), scanning tunneling microscopy (STM), x-ray photoelectron spectroscopy (XPS), and low-energy ion scattering (LEIS), have been characterized. We also calculated the surface energies of several possible surface terminations using density functional theory (DFT) to investigate surface stability.

## II. METHODS

### A. Experimental details

The single crystalline Ni₂In₃ sample was grown using the Czochralski technique [21]. The sample was cut along (001) surface into slices by a wire saw using a 50-µm-thick tungsten wire and boron carbide powder in a glycerol suspension to minimize surface damage of the crystal and material loss. Then, the surface was polished with a diamond paste of decreasing grain size down to 0.25 micrometers. The surface was prepared in a UHV chamber by standard procedures involving several cycles of a noble gas ion (Ar⁺) sputtering for 30 minutes, with an average drain current of ~6 µA and subsequently annealing for 1 to 3 hours between 300°C and 560°C. The temperature was measured by an optical pyrometer with an emissivity of 0.40 for the XPS, STM, and LEED experiments. As the emissivity for Ni-In was unknown, we used the respective value for Ni-Cu [22]. This procedure was repeated until LEED patterns were observed from the surface. The surface cleanliness was also checked after the preparation by XPS measurements, confirming the absence of contamination.

Then, the surface was freshly prepared and transferred to the STM chamber. We note that the preparation chamber

is equipped with LEED and interconnected to the XPS and STM chambers through the vacuum, ensuring that the surface conditions remain unchanged during transfer. STM images were obtained at room temperature using an Omicron Variable Temperature STM in constant current mode. The STM images presented in this paper were acquired with a tunneling current of 0.645 nA and a bias voltage of 0.14–1.7 V applied to the surface.

The LEIS experiment was conducted in a separate chamber. The surface was cleaned using the same method described above. For this experiment, the surface temperature was measured using a thermocouple in contact with the sample surface. LEIS spectra were obtained across a 2 mm × 2 mm area using a 3-keV He⁺ beam (ion dose ~1 × 10¹⁴ ions cm⁻²); only He⁺ ions with scattering energies in the range 1800–2800 eV were accumulated throughout the measurement to ensure a good signal-to-noise ratio for the Ni (~2275 eV) and In (~2575 eV) surface peaks. A commercial LEIS spectrometer, Qtac¹⁰⁰ (IonTOF), was used for all measurements, and analysis was performed in SurfaceLab 7 (IonTOF).

XPS data were collected at Institut Jean Lamour using a nonmonochromatic Al Kα source. Oxidation was carried out in UHV by introducing 99.999% pure oxygen gas through a leak valve. The partial pressure was maintained at about 1.3 × 10⁻⁶ mbar and the exposure was controlled by reaction time. The exposure was measured in units of Langmuir (L) (1L = 1.3 × 10⁻⁶ mbar s). Molecular hydrogen (99.99% pure) was dosed on the freshly prepared surface at room temperature. For the air exposure experiment, the surface was cleaned using the process outlined above, removed from the UHV chamber, and exposed to the atmosphere for two days before XPS measurements were performed.

### B. Computational details

Electronic structure calculations were performed using the plane wave Vienna ab initio simulation package (VASP) [23–26]. We applied the projector-augmented wave (PAW) method [27,28], within the spin-polarized generalized gradient approximation (GGA-PBE) [29,30]. We considered atomic valences to be 5s²5p¹ (In), 3d⁹4s¹ (Ni). Total energies were minimized until the energy differences became less than 10⁻⁶ eV between two electronic cycles during the structural optimizations. Atomic structures were relaxed till the Hellmann-Feynman forces were as low as 0.02 eV/Å.

Bulk calculations were performed using a 450 eV cut-off energy (Ecut) and a Γ-centered 9 × 9 × 5 k-point grid. Cell parameters have been optimized, and are found in good agreement with the literature: acut = 4.46 Å, ccut = 5.31 Å (aexp = 4.387 Å and cexp = 5.295 Å) [31]. In this paper, we use these experimental parameters to investigate the stability of several surface models. The formation entopy (−0.172 eV/atom) is also calculated and is in good agreement with the values in the literature (−0.16 eV/atom) [32].

## III. RESULTS AND DISCUSSION

Before discussing the results, we present the crystal structure of Ni₂In₃. It is an intermetallic compound, with a structure identical to AlNi-type, belonging to the trigonal
<figure><img src="image_3.png" /><figcaption>FIG. 1. (a) The unit cell of Ni₂In₃. (b) The side view of the (001) surface, illustrating a stacking of five atomic layers within the unit cell, marked as 0 to 4. The separations of specific planes are indicated. (c) Structures of the atomic models projected along [001], where the outermost atomic layer is selected at five distinct layers, 0 to 4, marked in (b). High-symmetry directions are indicated.</figcaption></figure>

space group of P-3m1 [31]. The lattice parameters are *a* = 4.387 Å and *c* = 5.295 Å, α = β = 90° and γ = 120° [31]. The trigonal unit cell of Ni₂In₃ is shown in Fig. 1(a), where each Ni atom is surrounded by four In atoms isolating Ni atoms from each other. Figure 1(b) shows the side view of the (001) surface. Within a unit cell, the crystal structure can be viewed as a stacking of five atomic planes, marked 0 to 4. The separation between planes 0 and 3 is 3.417 Å, and between planes 3 and the next 0 is 1.878 Å, with the sum of these distances equal to *c* (=5.295 Å). A comparison of these interplanar spacings with the step height observed in STM has helped to identify the bulk planes that terminate the surface.

Figure 1(c) shows atomic model structures projected along [001], where the outermost atomic layer is selected at five distinct layers, 0 to 4, marked in Fig. 1(b). We refer to these structures as surface models 0 to 4. We calculated the surface energies of these five model surfaces to investigate surface stability (see Sec. III E). Closely spaced planes can be considered as a single buckled plane, where In atoms are bonded with Ni atoms in the neighboring layer; for example, planes 0 and 1. Apart from surface model 2, all other models contain Ni in one of the top two layers. Model 2 contains only In atoms in the top two layers.

A. Surface chemical composition studied by XPS

XPS was used to investigate the influence of sputtering and annealing on surface chemical composition, and thus to optimize the surface preparation conditions (sample spectra will be presented in Sec. III F). The composition was calculated from the integrated intensity (area under the curve) of the Ni 2p$_{3/2}$ and the In 3d$_{5/2}$ peaks after removing the background using the Shirley function. The area was then normalized by the photoionization cross section [33], the inelastic mean free path of photoelectron [34], and the transmission function of

the analyzer [35]. The XPS data were analyzed using the CasaXPS program [35]. Figure 2 illustrates the change in surface chemical composition after sputtering and annealing at various temperatures in the range 300°–560°C.

After sputtering, the atomic percentage of In decreases to ~36% compared to 60% in the bulk, while that of Ni increases to ~64% from 40% expected from the bulk. This indicates that the sputtering process preferentially removes In, resulting in an excess of Ni at the surface. There are three

<figure><img src="image_4.png" /><figcaption>FIG. 2. The chemical composition of Ni₂In₃(001) surface after sputtering and annealing at various temperatures. The composition was determined using the integrated intensity of the Ni 2p$_{3/2}$ and the In 3d$_{5/2}$ peaks in XPS spectra, normalized by the photoionization cross section, the inelastic mean free path of photoelectron, and the transmission function of the analyzer. The bulk composition of In and Ni is marked by horizontal dotted lines.</figcaption></figure>
factors influencing the preferential sputtering of an element in intermetallic compounds or alloys: atomic weight, surface free energy, and bonding energy of the element. An element with lower weight, lower surface energy, and lower bonding energy is likely to be sputtered preferentially [36–38]. Since In is heavier than Ni, atomic weight is not responsible for preferential sputtering. However, In surface energy is lower than that of Ni. The surface free energy of the close-packed surface of body-centered tetragonal In is 0.488 J/m², while that of the close-packed surface of face-centered cubic Ni is 2.011 J/m² [39]. Therefore, we suggest that the lower surface-free energy of In causes preferential sputtering of this element.

After annealing the surface at 300°C, the atomic percentage of In increases to ~71%, and it remains steady as the sample temperature is raised to 560°C. In contrast, the atomic percentage of Ni decreases to ~29%, then remains stable with increased annealing temperatures. This change in atomic composition can be understood by the segregation of In to the surface during annealing. We measured XPS spectra from the annealed surface at grazing angles to understand the distribution of In in the surface region. The composition of In remains unchanged at a grazing angle of 70°, the maximum angle used in this experiment. This indicates that In is uniformly distributed in the surface region. If In were present only on the topmost layer, its composition would increase with the grazing angle.

To understand the surface segregation of an element in an alloy, several factors need to be considered. These factors include surface free energy, atomic radius, and bonding energy of the constituent elements [37]. The element with lower surface energy, larger radius, and lower binding energy is likely to segregate to the surface. The surface free energy of In is less than that of Ni, as stated above. Similarly, the atomic radius of elemental In is larger than Ni [40]. Therefore, we suggest that the segregation of In is owing to its lower surface free energy and the larger atomic radius of In compared to Ni.

## B. Low-energy electron diffraction

LEED patterns were taken from the surface after annealing at different temperatures from 300°C and 560°C, within an energy range of 10–250 eV, using a digital charge-coupled device camera with a computer-controlled data acquisition system. No LEED pattern emerged at a low annealing temperature of 300°C, indicating that this temperature is insufficient for the surface to become ordered. However, the surface began to exhibit a (1 × 1) LEED pattern displaying a hexagonal symmetry, after annealing at 310°C. The (1 × 1) pattern remains unchanged for annealing temperatures up to 560°C [Fig. 3(a)]. The lattice parameters determined from LEED patterns, by calibration with LEED patterns from the Ag(100) surface, are $a = b = 4.5 \pm 0.2$ Å, which are in good agreement with the expected values for the bulk terminated surface, 4.387 Å.

We found that the (1 × 1) pattern is unstable at a UHV pressure of 10⁻¹⁰ mbar. The surface forms a complex superstructure at this environment less than an hour after it is cleaned. Residual gases present in the chamber are a likely cause of this behavior. To confirm this further, we freshly prepared the surface after annealing at 560°C and fired the

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

FIG. 3. (a) LEED pattern from the freshly prepared Ni₂In₃(001) surface, showing a (1 × 1) structure. (b) LEED pattern from the same surface after firing the titanium sublimation pump, demonstrating the three domains of the c(2√3 × 4) rectangular superstructure. A section of the pattern is superimposed with the positions expected for the three domains. Both patterns were taken at 100 eV beam energy. (c) The positions of LEED spots expected for a single domain c(2√3 × 4)rect superstructure, where solid circles represent the position of the superstructure spots, while open circles denote the position of (1 × 1) spots. (d) The positions of the three domain c(2√3 × 4)rect LEED pattern, where the domain are rotated by 120° from each other. A yellow rhombus represents the unit cell of the (1 × 1) structure, whereas that of the c(2√3 × 4)rect domains is marked by red rectangles. The patterns in (c) and (d) were produced by LEEDpat [41].

titanium sublimation pump (TSP). The TSP upon operation releases the gases adsorbed on it, increasing the pressure in the chamber up to 10⁻⁵ mbar for several minutes. The surface adopts the identical superstructure immediately after the TSP is operated.

Figure 3(b) shows a LEED pattern after firing the TSP, which can be described as a c(2√3 × 4) rectangular superstructure with three different domains rotated by 120° from each other. The positions of LEED spots expected for a single domain and triple domains c(2√3 × 4)rect structure are shown in Figs. 1(c) and 1(d), illustrating the relationship between the superstructure and the substrate. The (1 × 1) LEED pattern can be recovered by annealing the sample at around 200°C.

One of the major contaminants inside the UHV chamber is hydrogen. Similarly, the gas released from the TSP is mostly hydrogen. It is therefore probable that hydrogen released from the TSP or present in the chamber causes the superstructure. We conducted XPS measurements on the surface after firing the TSP. XPS shows the presence of only In and Ni on the surface, ruling out the presence of any other gases on the surface. Noting that XPS is unable to detect hydrogen, we
<figure><img src="image_6.png" /><figcaption>FIG. 4. (a) STM image taken from the Ni₂In₃(001) surface after several cycles of sputtering and annealing at 550°C (400 Å × 400 Å, +0.866 V, 0.645 nA), demonstrating steps of different heights. (b) Line scan along a blue line in (a) crossing the steps. Steps heights are shown. (c) A high-resolution image of a terrace (240 Å × 240 Å, +1.11 V, 0.645 nA). (d) A magnified view of (c) with a unit cell of c(2√3 × 4)rect indicated by a red rectangle. The image is rotated to match the model in (e). (e) A model structure depicting a high-resolution STM image, where red spheres represent bright protrusions observed in STM, whereas the grid represents the unit cells of the (1 × 1) structure. A c(2√3 × 4)rect unit cell is shown by a red rectangle.</figcaption></figure>

conclude that the superstructure is caused by hydrogen, not by other adsorbates.

To confirm that hydrogen causes the superstructure, we freshly prepared the surface, which demonstrated the (1 × 1) structure, subsequently dosed 30 L of hydrogen, and then performed a LEED measurement. This hydrogen dosing indeed reproduces the c(2√3 × 4)rect superstructure, agreeing with the result after firing the TSP as described above. Annealing the surface after hydrogen-dosed recovers the (1 × 1) LEED pattern.

### C. Scanning tunneling microscopy

Annealing the surface from 420°C to 560°C for one hour yields a step-terrace morphology. However, to obtain large terraces the surface needed to be annealed longer, typically 2–3 hours above 530°C. An example of an STM image obtained after annealing at 560°C for two hours is shown in Fig. 4(a). The height of the steps is determined from a histogram of height distribution in the vicinity of the steps, and cross-checked by line scans. We analyzed over 100 steps from STM images taken from different parts of the surface to improve the statistical reliability. The most frequently observed step height is 5.2 ± 0.3 Å or a multiple of this height. The 5.2 Å high steps occasionally split into two smaller steps of 3.5 ± 0.2 Å and 1.7 ± 0.3 heights [see Fig. 4(a)].

These step heights are demonstrated in Fig. 4(b), which shows a height distribution along the direction indicated by a blue arrow in Fig. 4(a). The observed step heights are consistent with the separation of the model planes 0 and 3 [Fig. 1(b)], suggesting that the surface terminates at these

planes. This surface termination is further supported by surface energy calculations (Sec. III E).

It was not feasible to detect the (1 × 1) structure by STM owing to rapid contamination of the surface, as explained above. The surface undergoes reconstruction during sample transfer from the preparation chamber to the STM chamber. As anticipated from the LEED results, three domains producing the c(2√3 × 4)rect superstructure were observed in STM. However, STM reveals that the three domains do not coexist on the same terrace; they are observed on different terraces. Figures 4(c) and 4(d) show an example of a single domain. Brighter features form the central rectangular unit cells with dimensions of 2√3a and 4a, confirming the c(2√3 × 4)rect superstructure [also refer to the lattice representation in Fig. 4(e)]. Similar domains but rotated by 120° and 240° to the given domain were observed in different parts of the surface (see Fig. S1 within the Supplemental Material [42]). We typically select a 20 nm × 20 nm area of the surface to obtain high-resolution STM images, and observe only one domain within this area. Therefore, the minimum size of the domain is 20 nm × 20 nm. The root mean square roughness of the high-resolution STM images is 0.1 Å, which is comparable to that of metal surfaces.

We checked the quality of the sample by recording back reflection Laue x-ray diffraction from different parts of the surface. The vast majority of the sample produces a Laue pattern expected for a single crystal Ni₂In₃ with (001) orientation. However, a tiny section of the sample shows polycrystallinity of Ni₂In₃.

Moreover, the surface becomes hazy after annealing, indicating increased roughness, and becomes difficult to image with STM. Scanning electron microscopy (SEM) images recorded after annealing show an inhomogeneous surface morphology, with different contrasts. The chemical composition of these contrasting regions, analyzed using energy-dispersive x-ray spectroscopy (EDX), also varies. This suggests that the segregated In is unevenly distributed across the surface. However, the observed (1 × 1) LEED pattern indicates that the segregated In does not affect the overall surface symmetry.

### D. Low-energy ion scattering

We performed LEIS on the surface after sputtering and annealing cycles. LEIS probes the elemental composition of the outermost atomic layer of the crystal [43], in contrast to XPS, which provides the chemical composition of several atomic layers. For sputtering, a 3 mm × 3 mm area of the surface was exposed to a 1 keV Ar⁺ beam. Initially, the sample was exposed to the sputter beam until the LEIS spectra became constant with increasing sputter dose, hence removing surface layers to expose the bulk crystal. The sample was then treated similarly to our XPS experiments: it was heated to 400°C for one hour, cooled to room temperature over two hours, and then exposed to the sputter beam for 30 minutes (total dose = 3.0 × 10¹⁵ Ar⁺ ions). The sample was then heated to 400°C over a period of 15 minutes, and LEIS spectra were recorded periodically while the sample was held at 400°C for two hours. Figure 5 compares LEIS spectra obtained during this last sputtering and annealing cycle. The peak at ~2275 eV
<figure><img src="image_7.png" /><figcaption>FIG. 5. LEIS spectra recorded from the Ni₂In₃(001) surface after sputtering and annealing at 400°C for two hours using a 3-keV He⁺ beam. The energy ranges for each surface peak highlighted correspond to the upper and lower limits used for the baseline subtraction (refer to the text).</figcaption></figure>

corresponds to a signal from Ni, while the peak at ~2575 eV corresponds to a signal from In.

We also studied the change in the intensity of In and Ni as a function of sputter ion dose (depth profile) and annealing temperature. For this, the surface was freshly prepared by sputtering and annealing at 400°C for one hour and then cooled to room temperature. Subsequently, LEIS spectra were measured after varying ion doses, produc-

ing an elemental depth profile, and annealing at different temperatures up to 480 °C. For each annealing temperature, the sample was annealed for one hour, and a LEIS was recorded. The sample was then annealed at progressively higher temperatures without being cooled between steps. Figures 6(a) and 6(b) show the changes in intensity after these treatments. To calculate the intensity, the baseline was subtracted by a two-point linear subtraction over the energy ranges highlighted in Fig. 5, and the remaining signal, assigned to scattering from surface Ni and In, is integrated.

As seen in Figs. 6(a) and 6(c), initially, sputtering results in a decrease in the intensity of In, while the Ni content increases. With increasing sputter dose, the intensities of both In and Ni plateau. This is consistent with the XPS results, which show preferential sputtering of In. The LEIS results suggest that a thin (several nm), In-rich layer is present at the surface immediately following annealing. This is consistent with the hypothesis that In segregation occurs during heating, owing to the low surface free energy of In.

We note that the intensities of both In and Ni decrease slightly with annealing temperature from 295°C to 480°C. This is because the LEIS spectra were obtained *in situ*, and thus the LEIS signal is representative of the surface density, which may be expected to decrease with temperature owing to the expansion of the crystal. However, the ratio of In to Ni intensities follows the trend of increasing intensity with annealing temperature [refer to Fig. 6(d)]. We note that the surface undergoes reconstruction at this temperature, as observed by LEED and STM.

The LEIS spectra show well-defined In and Ni peaks, revealing that the outermost surface contains these elements

<figure><img src="image_8.png" /><figcaption>FIG. 6. The change in the integrated In and Ni intensities in LEIS from the Ni₂In₃(001) surface with sputtering at varying ion doses (a) and annealing at varying temperatures (b), along with the ratio of In to Ni intensities after the corresponding treatments (c), (d).</figcaption></figure>
after annealing. This suggests that the surface terminates at one of the surface models 0, 1, 3, or 4 [Fig. 1(c)]. For surface model 2, the In atoms in the topmost layer would block the incoming He⁺ ions from reaching the subsurface Ni. Hence, such a well-defined surface Ni peak would not be observed, unlike the one observed here [43]. Surface energy calculations rule out terminations for surface models 1 and 4 (Sec. III E). Therefore, a combination of LEIS and surface energy calculations suggests that the possible surface terminations are models 0 and 3.

## E. Surface energy calculations

Surface models were constructed by bulk truncation. After generating a 15-Å-thick slab with a (1 × 1) structure, we gradually removed atoms one by one from one side of the slab to systematically create distinct surface atomic arrangements. This process resulted in five surface models Fig. 1(c). Total energy calculations were performed to evaluate $T$, defined as $T = A(\gamma_{bottom} + \gamma_{top}^i)$, where A is the area of the surface cell, $\gamma_{bottom}$ is the surface energy of the bottom termination (the same for all slabs), and $\gamma_{top}^i$ is the surface energy of the $i$th surface model.

The quantity $T$ is given by the equation

$$T = E_{\text{slab}}^{i}(N_{\text{In}}, N_{\text{Ni}}) - N_{\text{In}}\mu_{\text{In}} - N_{\text{Ni}}\mu_{\text{Ni}},$$

where $E_{\text{slab}}^i(N_{\text{In}}, N_{\text{Ni}})$ is the total energy of the slab, $N_j$ ($j \in \{\text{In}, \text{Ni}\}$) is the number of atoms of type j in the slab, and $\mu_j$ is the chemical potential of the j-type element. The chemical potentials of elements j were evaluated in two different cases:

(1) in the Ni-rich limit, i.e., with $\mu_{\text{Ni}} = \mu_{\text{Ni}}^{\text{bulk}}$ and $\mu_{\text{In}} = \mu_{\text{In}}^{\text{bulk}} + \frac{\Delta H_f}{3}$,

(2) in the In-rich limit, i.e., with $\mu_{\text{In}} = \mu_{\text{In}}^{\text{bulk}}$ and $\mu_{\text{Ni}} = \mu_{\text{Ni}}^{\text{bulk}} + \frac{\Delta H_f}{2}$.

The chemical potentials $\mu_{\text{Ni}}^{\text{bulk}}$ and $\mu_{\text{In}}^{\text{bulk}}$ were calculated from the cohesive energy ($E_{\text{coh}}$) of elemental In and Ni. The calculated values [$E_{\text{coh}}(\text{Ni}) = 5.11 \text{ eV/atom}$ and $E_{\text{coh}}(\text{In}) = 2.34 \text{ eV/atom}$] are in good agreement with the literature [$E_{\text{coh}}(\text{Ni}) = 4.44 \text{ eV/atom}$ and $E_{\text{coh}}(\text{In}) = 2.52 \text{ eV/atom}$] [44].

The structural models with the lowest surface energies are models 0, 2, and 3 (Fig. 7). Models 1 and 4 are unlikely to be observed, as their surface energies are much higher ($\Delta T > 0.25 \text{ eV}$). As explained in Sec. III D, the LEIS results suggest that model 2 is less possible. If the terraces are formed on model planes 0 and 3, the expected step heights of the surface in STM are 3.417 Å and 1.878 Å [Fig. 1(b)]. These steps are observed in STM [Figs. 4(a) and 4(b)].

For the surface termination at models 0 and 3, the distance between two neighboring In atoms within the topmost layer is 4.387 Å, while the distance between a In atom in the topmost layer and an Ni atom in the layer below is 2.639 Å (for model 0) and 2.777 Å (for model 3). Therefore, In is likely to be bonded with Ni. This Ni-In bonding may be responsible for the stability of the crystal surface. Other binary intermetallic compounds, such as Zn-Pd, also demonstrate that surface terminations with both constituent elements are more stable than single-element terminations [13].

<figure><img src="image_9.png" /><figcaption>FIG. 7. Values of $T$ as a function of $\Delta\mu_{\text{In}}$ for the five models surfaces shown in Fig. 1(c). Refer to texts for the definition of $T$ and $\Delta\mu_{\text{In}}$.</figcaption></figure>

## F. Surface oxidation

Surface oxidation behavior was studied by XPS. The surface was dosed with various amounts of oxygen. First, 20 L oxygen was dosed on the freshly clean surface and XPS was measured. The surface was then recleaned, and exposed to 300, 600, and 900 L of oxygen, and XPS was applied after each dose. We also obtained XPS spectra from the surface after cleaning and then exposed it to air to compare with the results from the oxygen-exposed surfaces.

Figures 8(a)–8(c) display the In 3d, Ni 2p, and O 1s XPS peaks for the clean, oxygen-exposed, and air-exposed surfaces. As the oxygen exposure increases, a shoulder develops at higher binding energies across all three core levels. To better illustrate this shoulder formation, we plot the In 3d$_{5/2}$ and Ni 2p$_{3/2}$ peaks in Figs. 8(d) and 8(e), respectively, where the intensity of each spectrum has been normalized between 0 and 1. The shoulder is attributed to the formation of indium and nickel oxides. The peak and shoulder of O 1s are identified as pure oxygen, oxide, and Ni hydroxide, respectively. We will discuss the formation of oxides in detail later. After air exposure, the shoulder in In 3d$_{5/2}$ becomes the dominant peak [Fig. 8(d)]; while the shoulder in Ni 2p$_{3/2}$ remains largely unaffected [Fig. 8(e)]. The shoulder in O 1s also becomes the dominant peak after air exposure.

To determine the position and intensity of the peaks and shoulders, we fitted the spectra using a combination of asymmetric Gaussian and Lorentzian functions with a Shirley background. Figures 8(f)–8(i) provide examples of fitting for the In 3d$_{5/2}$ and Ni 2p$_{3/2}$ spectra from both the clean and oxygen-exposed surfaces. Here, we consider a combination of 60% Gaussian contributions and 40% Lorentzian contributions. All peak positions provided in this report were determined after fitting spectra with this approach. The error in peak position is $\pm 0.1 \text{ eV}$.

For the clean surface, the In 3d$_{5/2}$ and Ni 2p$_{3/2}$ can be fitted by a single peak located at 443.6 eV and 852.5 eV, as shown in Figs. 8(f) and 8(g), respectively. We attribute these peaks to intermetallic In and Ni. Since these peaks are close to
<figure><img src="image_10.png" /><figcaption>FIG. 8. (a)–(c) XPS spectra from clean (after annealing at 560 °C), oxygen-exposed and air-exposed (001) surface of Ni₂In₃ [(a) In 3d, (b) Ni 2p, and (c) O 1s]. Oxygen exposure was from 20–900 L. (d) (e) In 3d₅/₂ and Ni 2p₃/₂ spectra from (a) and (b), respectively, after normalizing minimum and maximum intensity between 0 and 1, illustrating the development of a shoulder at higher binding energy. (f) (g) In 3d₅/₂ and Ni 2p₃/₂ spectra from the clean surface after fitting with a single peak. The orange curve corresponds to the results of a fit. The method of fitting is described in the text. (h) (i) In 3d₅/₂ and Ni 2p₃/₂ spectra after 900 L oxygen exposure after fitting with two peaks: clean surface peaks (black) and oxide (blue) peaks. The dotted, dashed, and dash-dotted lines at 530.2 eV, 531.0 eV, and 531.6 eV in (c) correspond to the positions of the oxide, pure oxygen, and Ni hydroxide peaks, as reported in the literature ([45,46], and [47], respectively). The dotted lines in (c), (h), and (i) mark the positions of the oxide peaks, while the dashed lines in all the other spectra correspond to the positions of the core levels of the respective elements from the clean surface.</figcaption></figure>

the positions of the respective elemental In [48] and Ni [49], we are unable to confirm, with the given experimental energy resolution, whether there is charge transfer between In and Ni in the Ni₂In₃ intermetallic.

For the oxygen-exposed surface, the In 3d₅/₂ spectrum can be fitted by two peaks [Fig. 8(h)]. The large peak at 443.6 eV

is intermetallic In (pure In in Ni₃In₂). The small peak at 444.7 eV, which appears as a shoulder in the spectrum, is shifted to a higher binding energy by 1.1 eV. This shift is expected if In₂O₃ forms on the surface [48]. The oxide peak predominates over pure In on the air-exposed surface, suggesting that the air-exposed surface contains more In₂O₃ than pure
<figure><img src="image_11.png" /><figcaption>FIG. 9. The chemical composition of the Ni₂In₃(001) surface after annealing at 560°C and with increasing oxygen exposure (20-900 L). The atomic composition of the clean and air-exposed surfaces is also shown for comparison. The composition was determined from the In 3d₅/₂, Ni 2p₃/₂, O 1s, and C 1s XPS peaks.</figcaption></figure>

In, in contrast to the oxygen-exposed surface within the dosing regime used here (<1000 L), which contains more pure In.

Like In 3d₅/₂, Ni 2p₃/₂ after oxygen exposure can be fitted by two peaks: an intermetallic Ni peak located at 852.5 eV and a NiO peak located at 853.6 eV. The NiO peak is shifted to a higher binding energy by 1.1 eV from the pure Ni, which is consistent with the previous results [49]. In contrast to In, the intermetallic Ni peak remains dominant on the air-exposed surface. This suggests that Ni is less sensitive to oxidation than In. This observation also holds true for the oxygen exposure experiments. The intensity of the shoulder corresponding to In₂O₃ continuously increases with increasing oxygen exposure [Fig. 8(d)], whereas the shoulder corresponding to NiO saturates after 20 L of oxygen exposure [Fig. 8(e)]. After air exposure, an additional broad peak appears at 855.5 eV [indicated by an arrow in Fig. 8(e)], which is identified as a Ni(OH)₂ peak, by comparing the observed peak shift of 3.0 eV with a previous report [49].

We now present a qualitative analysis of the chemical composition of the surface after cleaning, oxygen exposure, and air exposure. The chemical composition was determined using the intensities of In 3d₅/₂, Ni 2p₃/₂, O 1s, and C 1s peaks after fitting (area under the curve) and the photoemission cross sections of the elements. Other factors such as instrumental response and inelastic mean free path of photoelectrons have been omitted. The composition of the surface after each treatment is shown in Fig. 9. As expected, only two elements, In and Ni, are detected on the clean surface. After oxygen exposure, pure oxygen and oxide peaks are observed in addition to pure In and Ni. However, air-exposed surfaces have hydroxide and carbon in addition to the species observed on the oxygen-exposed surface.

Pure In decreases from ~74% to ~27% from the clean surface to 900 L oxygen-exposed surface. In the same way, pure Ni decreases from ~28% to ~6% from the clean surface to the same exposure. The decrease in pure In and pure Ni is compensated by an increase in oxidized In (up to ~7% for 900 L) and oxidized Ni (up to ~2% for 900 L).

We note that oxygen was dosed on the c(2√3 × 4) surface. It was not possible to dose oxygen on the (1 × 1) surface owing to rapid contamination during the sample transfer from the preparation chamber to the XPS chamber, as explained above. However, we could measure LEED from the freshly prepared surface and subsequently from the surface after oxygen dosing. Interestingly, a (1 × 1) pattern was observed after dosing 300 L of oxygen (see Fig. S2 within the Supplemental Material [42]). However, it is unclear whether the (1 × 1) pattern arises from the clean substrate or an ordered oxygen (1 × 1) adsorbate phase. The LEED pattern disappeared after exposure to 900 L of oxygen, suggesting that this amount of oxygen destroys surface order.

## IV. SUMMARY

We have investigated the atomic structure, chemical composition, and oxidation behavior of the In₃Ni₂(001) surface using various techniques including low-energy electron diffraction (LEED), scanning tunneling microscopy (STM), x-ray photoelectron spectroscopy (XPS), and low-energy ion scattering (LEIS). The experimental results were complemented by density functional calculations of surface energies. The first step of this study was to optimize the surface preparation conditions using common techniques of sputtering and annealing. We found that sputtering the surface preferentially removes In, and annealing the surface promotes segregation of the same element into the surface from the bulk. The preferential sputtering and surface segregation are attributed to the low surface free energy of In compared to Ni. The overall surface symmetry remains to be (1 × 1) as revealed by LEED. However, the (1 × 1) surface is unstable upon exposure to ambient hydrogen present in the experimental chamber or to hydrogen dosed from an external source; it develops a three-domain c(2√3 × 4) rectangular structure as evidenced by both STM and LEED. Upon exposure of oxygen to the surface in an ultrahigh vacuum, In- and Ni-oxides are formed on the surface. A combination of the step height distribution observed by STM, LEIS intensity distribution, and surface energy calculations reveal that the outermost surface layer contains In atoms bonded with Ni atoms. This suggests that the stability of the surface may be caused by the formation of an intermetallic bond. Furthermore, the presence of catalytically active Ni atoms on the topmost layer, which are isolated from each other, would be beneficial for the catalytic properties of this material.

## ACKNOWLEDGMENTS

This work was partially supported by the Young Scientist Exchange Scheme of the European Integrated Center for the Development of Metallic Alloys and Compounds (ECMetAC). LEIS measurements were performed at the
University of Liverpool Low Energy Ion Scattering facility supported by UKRI-EPSRC Grant No. EP/P001297/1 and the University of Liverpool, maintained and operated as a shared research facility by the Faculty of Science and Engineering. A.Y.A. is grateful to Imam Abdulrahman Bin Faisal University, Saudi Arabia, for funding her Ph.D. studies. The authors gratefully acknowledge the Scanning Electron Microscopy Shared Research Facility at the University of Liverpool, and

in particular, Matthew Bilton, for his support and assistance with this work.

DATA AVAILABILITY

The data that support the findings of this article are not publicly available. The data are available from the authors upon reasonable request.

[1] M. Armbrüster, R. Schlögl, and Y. Grin, *Sci. Technol. Adv. Mater.* **15**, 034803 (2014).

[2] A. Álvarez, A. Bansode, A. Urakawa, A. V. Bavykina, T. A. Wezendonk, M. Makkee, J. Gascon, and F. Kapteijn, *Chem. Rev.* **117**, 9804 (2017).

[3] C. Quilis, N. Mota, E. Millán, B. Pawelec, and R. M. Navarro Yerga, *ChemCatChem* **16**, e202301496 (2024).

[4] M. Armbrüster, M. Behrens, K. Föttinger, M. Friedrich, E. Gaudry, S. K. Matam, and H. R. Sharma, *Catal. Rev.* **55**, 289 (2013).

[5] M. Armbrüster, K. Kovnir, M. Friedrich, D. Teschner, G. Wowsnick, M. Hahne, P. Gille, L. Szentmiklósi, M. Feuerbacher, M. Heggen et al., *Nat. Mater.* **11**, 690 (2012).

[6] M. Armbrüster, K. Kovnir, M. Behrens, D. Teschner, Y. Grin, and R. Schlögl, *J. Am. Chem. Soc.* **132**, 14745 (2010).

[7] K. Nozawa, N. Endo, S. Kameoka, A. P. Tsai, and Y. Ishii, *J. Phys. Soc. Jpn.* **80**, 064801 (2011).

[8] W. M. H. Sachtler, *Catal. Rev.* **14**, 193 (1976).

[9] J. Osswald, K. Kovnir, M. Armbrüster, R. Giedigkeit, R. E. Jentoft, U. Wild, Y. Grin, and R. Schlögl, *J. Catal.* **258**, 219 (2008).

[10] K. Kovnir, M. Armbrüster, D. Teschner, T. Venkov, L. Szentmiklósi, F. Jentoft, A. Knop-Gericke, Y. Grin, and R. Schlögl, *Surf. Sci.* **603**, 1784 (2009).

[11] W. Liu, L. Li, S. Lin, Y. Luo, Z. Bao, Y. Mao, K. Li, D. Wu, and H. Peng, *J. Energy Chem.* **65**, 34 (2022).

[12] A. P. Tsai, S. Kameoka, and Y. Ishii, *J. Phys. Soc. Jpn.* **73**, 3270 (2004).

[13] M. Lowe, A. Al-Mahboob, D. Ivarsson, M. Armbrüster, J. Ardini, G. Held, F. Maccherozzi, A. Bayer, V. Fournée, J. Ledieu, J. T. Sadowski, R. McGrath, and H. R. Sharma, *Phys. Rev. Mater.* **8**, 105801 (2024).

[14] C. Li, Y. Chen, S. Zhang, S. Xu, J. Zhou, F. Wang, M. Wei, D. G. Evans, and X. Duan, *Chem. Mater.* **25**, 3888 (2013).

[15] S. Dey and N. Mehta, *Chem. Eng. J. Adv.* **1**, 100008 (2020).

[16] H. Liu and D. He, *Catal. Surv. Asia* **16**, 53 (2012).

[17] V. A. Tsipouriari, Z. Zhang, and X. Verykios, *J. Catal.* **179**, 283 (1998).

[18] D. S. Costa, R. S. Gomes, C. Rodella, R. B. da Silva, R. Fréty, É. T. Neto, and S. T. Brandão, *Catal. Today* **344**, 15 (2020).

[19] L. Wang, S. Yamamoto, K. Hayashizaki, S. Nagamatsu, S. Malwadkar, T. Sasaki, and Y. Iwasawa, *ChemCatChem* **7**, 3248 (2015).

[20] D. Goodman, *Surf. Rev. Lett.* **02**, 9 (1995).

[21] P. Gille, in *Crystal Growth of Intermetallics*, edited by P. Gille (Walter de Gruyter, Berlin/Boston, 2019), pp. 61-90.

[22] https://ennologic.com/wp-content/uploads/2018/07/Ultimate-Emissivity-Table.pdf.

[23] G. Kresse and J. Hafner, *Phys. Rev. B* **47**, 558 (1993).

[24] G. Kresse and J. Hafner, *Phys. Rev. B* **49**, 14251 (1994).

[25] G. Kresse and J. Furthmüller, *Phys. Rev. B* **54**, 11169 (1996).

[26] G. Kresse and J. Furthmüller, *Comput. Mater. Sci.* **6**, 15 (1996).

[27] P. E. Blöchl, *Phys. Rev. B* **50**, 17953 (1994).

[28] G. Kresse and D. Joubert, *Phys. Rev. B* **59**, 1758 (1999).

[29] J. P. Perdew, K. Burke, and M. Ernzerhof, *Phys. Rev. Lett.* **77**, 3865 (1996).

[30] J. P. Perdew, K. Burke, and M. Ernzerhof, *Phys. Rev. Lett.* **78**, 1396(E) (1997).

[31] E. Hellner, *International Journal of Materials Research* **41**, 401 (1950).

[32] C. E. D. Toro, S. R. de Debiaggi, and A. M. Monti, *Phys. B: Condens. Matter* **407**, 3236 (2012).

[33] J. Scofield, *J. Electron Spectrosc. Relat. Phenom.* **8**, 129 (1976).

[34] M. P. Seah and W. A. Dench, *Surf. Interface Anal.* **1**, 2 (1979).

[35] N. Fairley, Casa Software Ltd, Software Package for the Analysis of XPS Results, CasaXPS version 2.3.17dev6.6o, http://www.casaxps.com.

[36] L. C. Feldman and J. W. Mayer, *Fundamentals of Surface and Thin Film Analysis* (Elsevier Science, Amsterdam, 1986).

[37] C. D. Yuen, B. Unal, D. Jing, and P. A. Thiel, *Philos. Mag.* **91**, 2879 (2011).

[38] P. Sigmund, *Phys. Rev.* **184**, 383 (1969).

[39] L. Vitos, A. V. Ruban, H. L. Skriver, and J. Kollár, *Surf. Sci.* **411**, 186 (1998).

[40] T. O. Owolabi, K. O. Akande, and S. O. Olatunji, *Int. J. Adv. Inf. Sci. Technol.* **3**, 105 (2014).

[41] K. E. Hermann (FHI) and M. A. Van Hove (HKBU), http://www.fhi-berlin.mpg.de/KHsoftware/LEEDpat/index.html (2014).

[42] See Supplemental Material at http://link.aps.org/supplemental/10.1103/PhysRevMaterials.9.035801 for containing two figures. Figure S1 shows STM images of the Ni₂In₃(001) surface with three domains of the c(2√3 × 4)rect structure, each rotated by 120° relative to one another. Figure S2 displays LEED patterns from the freshly prepared Ni₂In₃(001) surface and after exposure to 300 L of oxygen.

[43] C. V. Cushman, P. Brüner, J. Zakel, G. H. Major, B. M. Lunt, N. J. Smith, T. Grehl, and M. R. Linford, *Anal. Methods* **8**, 3419 (2016).

[44] C. Kittel, *Introduction to Solid State Physics*, 7th ed. (John Wiley & Sons, Hoboken, NJ, 1996).
[45] P. J. Nugent, G. Simutis, V. R. Dhanak, R. McGrath, M. Shimoda, C. Cui, A. P. Tsai, and H. R. Sharma, Phys. Rev. B 82, 014201 (2010).

[46] J. F. Moulder and J. Chastain, Handbook of X-ray Photoelectron Spectroscopy: A Reference Book of Standard Spectra for Identification and Interpretation of XPS Data (Perkin-Elmer Corporation, Eden Prairie, Minnesota, USA, 1992).

[47] V. Crist, https://xpslibrary.com/summaries-of-be-tables-nist/ (2018).

[48] Z. M. Detweiler, S. M. Wulfsberg, M. G. Frith, A. B. Bocarsly, and S. L. Bernasek, Surf. Sci. 648, 188 (2016).

[49] R. P. Furstenau, G. McDougall, and M. A. Langell, Surf. Sci. 150, 55 (1985).