Electronic properties of the selenium passivated
GaP(111)B surface: Towards growth of large scale
quasi-van der Waals 2D/3D heterostructure

Niels Chapuis, Corentin Sthioul, Aymen Mahmoudi, Meryem Bouaziz, Christophe Coinon, Louis Thomas, Davide Romanin, Gilles Patriarche, Fabrice Oehler, Abdelkarim Ouerghi, et al.

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Niels Chapuis, Corentin Sthioul, Aymen Mahmoudi, Meryem Bouaziz, Christophe Coinon, et al. Electronic properties of the selenium passivated GaP(111)B surface: Towards growth of large scale quasi-van der Waals 2D/3D heterostructure. Physical Review Materials, 2025, 9 (7), pp.074002. 10.1103/pv8b-89gg. hal-05165772

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Electronic properties of the selenium passivated $GaP(111)_B$ surface:
towards growth of large scale quasi-van der Waals 2D/3D heterostructure

Niels Chapuis¹, Corentin Sthioul¹, Aymen Mahmoudi², Meryem Bouaziz², Christophe Coinon¹, Louis Thomas¹, Davide Romanin², Gilles Patriarche², Fabrice Oehler², Abdelkarim Ouerghi² and Xavier Wallart¹

¹Univ. Lille, CNRS, Centrale Lille, Junia, Univ. Polytechnique Hauts-de-France, UMR 8520 - IEMN - Institut d'Electronique de Microélectronique et de Nanotechnologie, F-59000 Lille, France

²Université Paris-Saclay, CNRS, Centre de Nanosciences et de Nanotechnologies, 91120, Palaiseau, Paris,
France

Corresponding author: niels.chapuis@c2n.upsaclay.fr

**ABSTRACT.** In recent years, several well-established 3D substrates have been considered for the Van der Waals epitaxy of 2D-Transition Metal Dichalcogenides, yet requiring a proper surface passivation treatment. In this work, we study the structural and electronic properties of well-ordered Se-terminated $GaP(111)_B$ surfaces by molecular beam epitaxy. We emphasize the key role of using a combination of atomic Hydrogen and cracked phosphine at a relative low temperature to prepare a smooth P-rich surface with a (2 × 2) reconstruction of n and p doped $GaP(111)$. We then investigate the evolution of the surface electronic properties after a selenium flux exposition. The $GaP(111)_B$ surface exhibits a systematic change in reconstruction from a (2 × 2) to a (1 × 1) pattern, while showing a change in chemical environment mainly in the outermost P-surface layer. We evidence the absence of a GaSe layer formation through angle resolved photoemission spectroscopy measurements, while demonstrating that the band structure of the Se-passivated surface originates from the $GaP(111)_B$ one. We then emphasize the predominant effect of a low selenium flux, while demonstrating an effective partial electronic passivation of the surface through the evolution of the band bending and the surface dipole measured with ultra violet photoemission spectroscopy. This realization of Se-terminated $GaP(111)_B$ substrate paves the way to new 2D/3D heterostructures development with great interests in nanoelectronic and optoelectronic applications.

**Keywords:** 2D materials, MBE, Electronic Band Structure, Van der Waals

# I. INTRODUCTION

Within the growth of heterostructures, van der Waals (vdW) epitaxy for which the substrate and epilayer do not form covalent bonds but are only linked by vdW interactions is very attractive since it relaxes the stringent lattice matching condition encountered in standard epitaxy [1]. Ideally, it requires a vdW substrate, i.e. a substrate without any surface dangling bonds, a condition fulfilled by two-dimensional (2D) materials. However, substrates made of 2D materials and compatible with microelectronic applications and epitaxial techniques are relatively scarce with the noticeable exception of graphene, which is metallic (and not a natural substrate), hexagonal boron nitride and micas. The search for alternative commercially available substrates, more compatible with microelectronics and epitaxy requirements is then mandatory. In this perspective, previous work showed that Se-terminated
GaAs(111) surfaces are well suited for vdW epitaxy of lamellar materials [2–8]. Besides, in the last two decades, numerous scientific studies focusing on 2D materials and especially on transition metal dichalcogenides (TMDs) [9] have demonstrated that the crystalline quality improves at high growth temperatures [10,11]. That is why, besides GaAs, GaP with a higher thermal stability than GaAs [12] deserves consideration [13]. Moreover, GaP has a lattice parameter (5.45 Å) close to that of silicon (5.43 Å) which could promote future integration of TMDs on silicon if proper GaP/Si(111) templates are developed.

In the last two decades, following the discovery of graphene, numerous scientific studies have focused on 2D materials and especially on transition metal dichalcogenides (TMDs) [9] with a main concern about material crystalline quality which improves at high growth temperatures [10,11]. That is why, besides GaAs, GaP with a higher thermal stability than GaAs [12] deserves consideration [13].

However, the preparation and properties of the GaP(111)B surface have been far less studied than the GaAs one for instance and, to our knowledge, no experimental work on its Se passivation has yet been reported. Regarding the surface cleaning, using ion bombardment followed by an annealing around 500°C in Ultra-High Vacuum (UHV) conditions led to either a (1 × 1) Low Energy Electron Diffraction (LEED) pattern [14–16] or to an incomplete (17 × 17) one [17,18], this latter turning to a 1 × 1 one around 250°C. Using LEED and Scanning Tunnelling Microscopy (STM), Hattori et al. [19] further discussed and interpreted this (17 × 17) reconstruction as six equivalent mirror-symmetric domains, each domain having a stripe structure. More recently, P. Kleinschmidt et al. [20] showed by STM that the GaP(111)B surface prepared in a Metal Organic Vapor Phase Epitaxy (MOVPE) environment with annealing under H₂ exhibits domains with different reconstructions with no long range order and an overall (1 × 1) LEED pattern. They explained these domains by Density Functional Theory (DFT) calculations taking into account the hydrogen-rich environment of MOVPE. In UHV environment and under phosphorus-rich conditions, these calculations predict a stable (2 × 2) P-terminated reconstruction. During homoepitaxy by Gas Source Molecular Beam Epitaxy, Barakat et al. [21] described different Reflection High Energy Diffraction (RHEED) patterns from (2 × 2) to (3 × 3) along the different growth stages, the resulting surface morphology being rather rough whatever the growth conditions. Finally, if a rather high temperature is required to get a deoxidized surface without Ar⁺ bombardment, Ga droplets are prompt to form beyond 700°C [22]. Regarding the electronic properties, using Electron Energy Loss Spectroscopy, Jakobi [23] deduced the existence of filled surface states 0.8 eV above the valence band edge (VBE) for a P-rich surface and of empty ones 1.4 eV above the VBE for a Ga-rich surface. On the other hand, Dai et al. [14] from Total Current Spectroscopy revealed intrinsic surface states 0.4 eV below the VBE associated to the surface P atoms and extrinsic ones at 1.1 eV above the VBE originated from defects caused by the Ga enrichment of the surface during its preparation.

The evolution of the GaP surface properties upon chalcogenide treatment has been scarcely reported with few studies on the sulphur-treated surfaces [24]. For GaP(001) surfaces exposed to H₂S vapour, Fukuda et al. [25] only detected Ga-S bonds as well as Lu et al. [26] using a (NH₄)₂S solution, with the S atom occupying a bridge site on the top surface [25]. On the H₂S treated GaP(110) surface [27], P-S bonds are also detected but they are weaker than the Ga-S ones. More recently, on (NH₄)₂S treated GaP(001), GaP(111)A and GaP(111)B surfaces, Suzuki et al. [28] evidenced three components in the sulphur X-ray Photoelectron Spectroscopy (XPS) spectrum associated to Ga-S bonds at the surface, Ga-S at the subsurface and P-S bonds. The latter component nearly vanishes after a 400°C annealing contrarily to the two other ones whereas a flat band situation is recovered for the (001) surface. Upon annealing at higher temperature, Itagaki et al. showed that sulphur desorbs at temperature of about 550°C
from $(NH_4)_2S$ prepared $GaP(111)_A$ surfaces [29]. According to DFT calculations, the sulphur or selenium passivated $GaP(001)$ surfaces exhibit similar behaviour. The most stable structure for the induced (2 × 1) reconstruction corresponds to S or Se surface atoms occupying bridge sites and P atoms in the third layer replaced by chalcogen ones [30]. Finally, under selenium partial pressures in the $10^{-2}-10^{-1}$ T range, Bezrvadin et al. reported the formation of a $Ga_2Se_3$ surface phase [31]. But a detailed investigation of the $GaP(111)_B$ surface under selenium treatment under an ultra-high vacuum environment is still missing.

In this study, we present results relative to the preparation and characterization of well-ordered $GaP(111)_B$ surfaces in a Molecular Beam Epitaxy (MBE) environment. We show that avoiding too high substrate temperature during the deoxidation step allows clean and smooth P-rich surfaces exhibiting a (2 × 2) reconstruction to be prepared. We then investigate the surface evolution upon exposure to a selenium flux and discuss the stability as well as the electronic properties of the resulting $GaP(111)_B-Se$ surface.

## II. METHODS

2-inch $GaP(111)_B$ substrates were purchased from Ensemble3 ltd, Warsaw with two different doping: n+(2.1-3.6×10¹⁸/cm³, sulphur doped) and p (6-7.1×10¹⁷/cm³, zinc doped). The clean surfaces were prepared in a III-V MBE reactor Riber Compact 21 fitted with standard element-III effusion cells, a high temperature injector and a RF plasma source to produce a $P_2$ flux through phosphine ($PH_3$) cracking and an atomic hydrogen (AH) flux for substrate cleaning, respectively. After transfer under UHV, the samples were exposed to a selenium flux in a second MBE reactor from Vinci Technologies thanks to a selenium valve cracker cell. Both MBE reactors are coupled under UHV to a surface analysis chamber equipped with a Physical Electronics 5600 system allowing XPS, Ultraviolet photoelectron spectroscopy (UPS) and LEED characterizations. XPS was performed with a monochromatic Al Kα anode (1486.6 eV) making a 90° angle with the hemispheric photoelectron analyzer whose acceptance angle was set to 14° whereas the electron take-off angle with respect to the surface plane can be varied from 25° to 75°. Core level decomposition is achieved using pseudo-Voigt functions within the PHI MultiPak™ software [31]. The branching ratios are theoretical ones and the splitting energies are fixed at 0.86 eV between Se $3d_{5/2}$ and Se $3d_{3/2}$, 0.84 eV between P $2p_{3/2}$ and P $2p_{1/2}$ and 0.45 eV Ga $3d_{5/2}$ and Ga $3d_{3/2}$. For X-ray photoelectron diffraction (XPD) experiments, a Mg Kα anode was used (1253.6 eV) and the analyzer acceptance angle reduced to 4°. In UPS, the valence band spectra were recorded around the Γ point of the Brillouin zone with the HeI radiation (21.2 eV) generated by a He plasma UV source. Survey spectra were acquired with a 58.7 eV pass energy, which was reduced to 12 eV for high-resolution spectra. The samples were then characterized *ex-situ* by Atomic Force microscopy (AFM), either with a Multimode or Dimension microscope working in the tapping mode. The ARPES experiments were performed at the CASSIOPEE beamline of the SOLEIL synchrotron light source, equipped with a Scienta R4000 hemispherical electron analyzer with a ±15° angular acceptance (Scienta Wide Angle Lens). The experiment was performed at T=70 K with a total angle and energy resolution of 0.25° and 16 meV, respectively [33]. To prevent sample oxidation during the transfer to the synchrotron, a 30 nm-thick amorphous selenium layer was deposited just after the sample preparation in the MBE chamber. This capping layer was then desorbed by a moderate heating up to 220°C in the preparation chamber of the CASSIOPEE beam line. DFT calculations for bulk $GaP(111)$ have been performed in the plane-wave basis expansion as implemented in the Quantum ESPRESSO package [34,35]. We employed an ultrasoft pseudopotential for P atoms, while we considered a projector-augmented-wave (PAW) pseudopotential for Ga atoms: the wave-function and electron density cut-offs have been fixed at 90 Ry and 720 Ry respectively. Electronic band structure and energetics have
been obtained through Perdew-Burke-Ernzerhof (PBE) exchange-correlation functional. The Brillouin zone has been sampled with a k-vector mesh of 8 × 8 × 4 points in the hexagonal cell of [111]GaP, assuring convergence of the total energy per atom. 2D Fermi surfaces have then been obtained through constant energy cuts of electronic eigenvalues at 0.5 eV below the top of the valence band, on a kₓ vs kᵧ grid of 60 × 60 at the center of the Brillouin zone (kz = 0.0).

### III. RESULTS AND DISCUSSION

We first discuss the preparation and the electronic properties of the GaP(111)B (2 × 2) surface before studying the effect of the selenium treatment on its structure and band structure.

#### A. Structural and electronic properties of GaP(111)B (2 × 2) deoxidized surface

We first tried to deoxidize the GaP(111)B surface by a thermal annealing under a P₂ flux monitored by RHEED. The initial faint (1 × 1) pattern reinforced progressively with increasing substrate temperature and phosphine flow rate up to 8sccm. Around 670°C, it became clearer without any reconstruction but spotty, indicative of deoxidation and of some surface roughening at the same time. The AFM image of the resulting surface (see Fig. S1(a)) revealed a rather rough surface with a RMS roughness of 0.4 nm.

Obviously, to avoid surface roughening and even Ga droplets formation induced by high substrate temperature annealing [36], it is preferable to make the deoxidation at lower temperatures. Since an atomic hydrogen flux has proven to be useful in the case of GaAs surfaces [37,38], we then performed the deoxidation using a combination of P₂ and AH fluxes in three steps as the substrate temperature was ramped up: AH only then AH + P₂ and finally P₂ only. The optimization of this procedure relied on AFM imaging for the surface morphology and on XPS measurements to track residual traces of oxygen on the surface. We noticed that a minimum exposure time of 10 min. at substrate temperatures above 250-300°C was required to get a XPS O 1s signal below the detection limit. After several trials, the optimized sequence was defined as: AH from 300 to 400°C, AH+ P₂ from 400 to 480°C and P₂ only from 480 to 540-550°C. The corresponding RHEED pattern evolved from a faint starting (1 × 1) to a well-defined (1 × 1) up during AH and AH + P₂ exposures and finally turned to a (2 × 2) one around 500°C under P₂ only. The corresponding RHEED and LEED patterns together with the AFM image where atomic steps are clearly resolved and a surface roughness of 0.15 nm are displayed in Figure 1. This finding is in good agreement with the calculations of Kleinschmidt et al. [20] predicting a (2 × 2) reconstruction with P ad-trimers on top of the P-terminated GaP(111)B surface for a GaP(111)B surface prepared under UHV in a phosphorus-rich environment. The P-trimers periodic arrangement in both surface directions is the origin of the (2 × 2) reconstruction. Upon increasing the substrate temperature, the (2 × 2) reconstruction progressively vanished and disappeared around 570-580°C under 8sccm PH₃ (540°C under 2 sccm PH₃) giving rise to a faint (√19 × √19) diagram around 590-600°C in a way rather similar to what was observed on GaAs(111)B [39,40]. Decreasing the substrate temperature led to the transition back to the (2 × 2) diagram around 540°C under 8 sccm PH₃. Nevertheless, AFM revealed that the morphology of surfaces being annealed above 550°C is characterized by dots that we associated with the formation of Ga droplets (see Fig. S1(b)). In conclusion, a combination of AH and P₂ fluxes allowed for a complete deoxidation of the GaP(111)B surface with a smooth morphology.

The normalized XPS Ga 3d and P 2p core level lines (CLs) recorded on the surface of a n-type sample prepared as discussed above are shown in Figures 1(e) and 1(f) for 2 electron takeoff angles with respect to the surface : 25°
(surface sensitive) and 75° (bulk sensitive). The slight binding energy shift observed on both CLs is associated to band bending effects due to the rather high doping level of the n-type substrate (2-3×10¹⁸/cm²). Whereas the Ga 3d CL does not show strong variations with the detection angle, surface components are clearly present on the P 2p CL. Figures 1(e) and 1(f) display the decomposition of both CLs for a takeoff angle of 25°. The Ga 3d one is composed of one major component associated to the Ga-P bonds at 19.98 eV. On the P 2p CL, besides the main doublet reflecting the P-Ga bonds at 129.63 eV, a clear surface component (SCL) appears on the high binding energy side (S1) at 130.34 eV and a second one (S2) on the low binding energy side at 128.88 eV. Two SCLs have already been evidenced on some reconstructions of the GaP and InP(001) surfaces [41] whereas a SCL comparable to S2 has been reported on GaP(111)B surfaces prepared under hydrogen [20]. Such surface components have been also observed on the GaAs(111)B (2 × 2) surface and interpreted in terms of the As trimer model [42–44]. Considering previous calculations predicting a (2 × 2) reconstruction based on P ad-trimers [20], we infer that the GaAs(111)B (2 × 2) surface structural model also holds for the GaP(111)B surface. Therefore, the high binding energy SCL S2 can be associated to P-P bonds formed within the trimers by the excess P atoms. The low binding energy SCL S1 might be related to the outermost P atoms on the surface bonded to three Ga atoms and carrying a filled dangling bond. The p-type samples exhibit similar behavior.

<figure><img src="image_2.png" /><figcaption>FIG 1. 2 × 2 RHEED pattern along <11-2> (a) and <1-10> (b) obtained after the optimized deoxidation process in three steps at a final temperature of 550°C, with white lines representing intensity profiles with global integration of the window. Associated LEED pattern recorded at 40 eV (c) and AFM image (d). Normalized Ga 3d CL (e) and P 2p CL (f) recorded for two take-off angles on a deoxidized n-type GaP(111)B surface with the associated Ga 3d and P 2p CLs decompositions for a take-off angle θ=25°. S1: P-P bonds at the surface; S2: P-Ga bonds at the surface.</figcaption></figure>

Figure S2(a) illustrates the valence band spectra recorded on n⁺ and p surfaces. The energy of the top of the valence band with respect to the Fermi level at the surface E<sub>F</sub> - E<sub>vsurf</sub> is determined by the intersection of a linear extrapolation of the low binding energy part of the valence band with the background (dashed lines in Fig. S2(a)). Taking into account the doping level of both samples and the data for the electronic structure of GaP [45] allows calculating the position of the valence band maximum in the bulk, E<sub>F</sub> - E<sub>vbulk</sub>. The differences between E<sub>F</sub> - E<sub>vsurf</sub> and E<sub>F</sub> - E<sub>vbulk</sub> is a measure of the upwards surface band bending and downwards one for the n⁺ and p samples respectively (see Table I). This behavior suggests the existence of two surface state bands within the bandgap
centered around 0.57 eV and 1.59 eV above the valence band maximum. These values are in rather good qualitative agreement with those reported by Jakobi et al. [23] who mentioned two surface states at 0.8 and 1.4 eV above the valence band maximum.

Finally, using a negative sample polarization allows the determination of the electron emission threshold on the high binding energy side of the valence band spectrum and therefore the work function $\Phi$ (see Fig. S2(b)). Considering the known electron affinity $\kappa$ of GaP (3.8 eV) [44] leads then to the determination of the surface dipole according to [46]:

$$ \Phi = E_g - (E_F - E_{vbulk}) + \kappa + \Phi_b + \Phi_s $$

where $E_g$ is the bandgap energy, $E_F$ the Fermi level, $E_{vbulk}$ the energy of the top of the valence band in the bulk, $\kappa$ the electron affinity, $\Phi_b$ the band bending and $\Phi_s$ the surface dipole. Table I summarizes the measured values and the deduced surface dipoles and Figure 2 gives a picture of the surface electronic structure for the two different doping types considered.

<figure><img src="image_3.png" /></figure>

<figure><img src="image_4.png" /><figcaption>FIG 2. Band structure schematic of deoxidized (a) n-type and (b) p-type GaP(111)$_B$ surface.</figcaption></figure>

<table><caption>TABLE I. Surface electronic properties for a deoxidized GaP(111)$_B$ (2 × 2) surface of a n+ and p-type substrate. Measured position of the Fermi level at the surface (E$_F$-E$_{vsurf}$), calculated position of the Fermi level in the bulk (E$_F$-E$_{vbulk}$), deduced band bending (Φ$_b$), measured work function (Φ) and deduced surface dipole (Φ$_s$).</caption><thead><tr><th>Sample</th><th>Carrier density (cm<sup>-3</sup>)</th><th>E<sub>F</sub>-E<sub>vsurf</sub> (eV)</th><th>E<sub>F</sub>-E<sub>vbulk</sub> (eV)</th><th>Φ<sub>b</sub> (eV)</th><th>Φ (eV)</th><th>Φ<sub>s</sub> (eV)</th></tr></thead><tbody><tr><td>n<sup>+</sup></td><td>n~2 - 3 × 10<sup>18</sup></td><td>1.59</td><td>2.2</td><td>0.61</td><td>4.31</td><td>-0.16</td></tr><tr><td>p</td><td>p~6 - 7 × 10<sup>17</sup></td><td>0.57</td><td>0.09</td><td>-0.48</td><td>5.12</td><td>-0.37</td></tr></tbody></table>

## B. Se-passivated GaP(111)$_B$ (2 × 2) surface: Evolution of the surface and the band structure upon Se exposure
After the clean GaP(111)B (2 × 2) surface was prepared according to the above procedure, the sample was transferred under UHV in a MBE chamber dedicated to Se-based materials. It was then annealed to 250°C under vacuum before being exposed to a selenium flux FSe while increasing the temperature to a maximum value TSe. The substrate remained at this temperature for 3 minutes before cooling down.

As soon as the GaP(111)B (2 × 2) surface was exposed to the Se flux at 250°C, the (2 × 2) RHEED and LEED patterns turned to a (1 × 1) one (Fig. 3(a) and 3(b)), remaining unchanged up to 700°C. Actually, this (1 × 1) pattern was stable up to more than 750°C, a strong degradation appearing around 800°C but in the following, we do not consider anymore the annealing temperatures above 700°C. For surfaces being annealed under selenium up to 700°C, AFM did not reveal any change of the surface morphology (Fig. 3(b) and 3(d)), with clear atomic steps of identical height (~3.5 Å), serrated shape and same mean roughness (~0.2 nm). On the contrary, the XPS survey spectrum of such surfaces exhibits the characteristic CLs of selenium besides those of gallium and phosphorus (Fig. 3(e)).

Figure 3 also displays the decomposition of the Se 3d, P 2p and Ga 3d CLs for a takeoff angle of 25° after Se exposure at 700°C. The P 2p CL (Fig. 3(f)) does not exhibit anymore any surface component S1 at high binding energy related to excess P atoms in the trimers. This is consistent with the change of reconstruction since the (2 × 2) one observed on the bare surface was due to the trimer periodic arrangement. Contrarily, the low binding energy SCL S2 associated to outermost P atoms bonded to three Ga ones still remains with a reduced intensity. Thus, we assume that selenium mainly reacts with the P-terminated surface, though some P-atoms remain in the outermost layer, probably giving rise to a mixed P-Se layer as discussed in the case of GaAs(111)B-Se [47,48].

The Ga 3d CL after Se-passivation (Fig. 3(g)) exhibits an additional component at higher binding energy (20.95 eV) that we ascribe to the formation of surface Ga-Se bonds as the Se electronegativity is larger than the P one. Finally, the Se 3d CL in Figure 3(h) is composed of two doublets Se1 and Se2 at binding energies of 55.38 eV and 54.38 eV, respectively. As proposed by Scimeca et al. in the case of the GaAs(111)B surface exposed to Se [46], we interpret Se1 as resulting from the Se-Ga bonds on the top surface whereas Se2 is related to some Se atoms present in the sub-surface layer.
<figure><img src="image_5.png" /><figcaption>FIG. 3. LEED pattern at 40 eV and 2 × 2 µm² AFM images of the GaP(111)B surface before (a,c) and after the Se treatment at 700°C (b,d). XPS Survey spectra after deoxidation and after Se treatment recorded with a take-off angle of θ=45° (e). Decompositions of P 2p (f), Ga 3d (g) and Se 3d (h) CLs of a n-type GaP(111)B substrate after a passivation at T<sub>Se</sub>=700°C and F<sub>Se</sub>=3.9×10<sup>-6</sup> Torr. The take-off angle is set at θ=25°.</figcaption></figure>

In order to study the electronic properties of the Se-terminated GaP(111), we performed band structure measurements by ARPES. In Figure 4(a) we report the Se-terminated GaP band structures along the ΓK direction in the Brillouin zone (BZ), obtained with a photon energy of 60 eV. The contribution from the GaP(111)B substrate is high and we clearly see the typical GaP manifold of light hole (LH), heavy hole (HH) and spin-orbit split (SO) bands. We note that the VBM of GaP is located at the Γ point, with a BE VBM GaP(111)B = -2 ± 0.01 eV. The measured band structure has been compared with the theoretical calculation along the ΓK direction of the bulk GaP(111) BZ in Figure 4(b), in the absence of any Se atom. The DFT calculation shows that bulk GaP(111) is a
compound semiconductor material with a direct bandgap in the hexagonal cell which is slightly smaller than the experimental values. From the excellent agreement between the measurements and the DFT calculations, we conclude that the band structure originates from the GaP(111)$_{B}$ without any change associated with Se.

For the sake of clarity, it is important to note that the value of the band gap extracted using DFT and semi-local functionals is a lower estimation of the experimental value. Our ARPES measurements show that the Fermi level lies within the gap. Despite the VBM being located at 2 eV, no conduction band states were observed. As such, we only compare the DFT valence bands with ARPES measurements.

Figure 4(c) shows an isoenergetic contour in the (kₓ, kᵧ) plane obtained at the valence band maximum. In particular, we observe the evidence of the flowering of the electronic band structure at the Γ-point of the BZ. To validate the experimental results, we performed density functional theory (DFT) calculations for bulk GaP oriented along the [111] crystallographic direction in its hexagonal unit cell, without any Se atom. Thus, at kₓ = 0, we obtain a flower-like constant energy surface at the VBM around the center of the Brillouin zone, which originates from three bulk valence bands folded at Γ due to the modification of the unit cell from hexagonal-[111] to cubic-[100]. The measured band structure is in good agreement with the theoretical calculation (Fig. 4(d)). This result confirms the absence of GaSe on GaP(111)$_{B}$ [49,50].

<figure><img src="image_6.png" /><figcaption>FIG. 4. (a) ARPES measurements of the Se terminated GaP(111) along the ΓK direction of the Brillouin zone measured at hv = 60 eV, (b) comparison between DFT electronic bands of pure bulk GaP(111) and ARPES measurements of the Se terminated GaP(111). The Fermi level position is at the zero binding energy, (c) iso-energy</figcaption></figure>
cuts of the first Brillouin at VBM-0.5 eV, (d) comparison between DFT Fermi surface of pure bulk GaP(111) and iso-energy cuts of the Se terminated GaP(111) at the G-point.

In order to confirm that Se atoms are mainly present in the top surface layer, we performed X-ray- photoemission diffraction experiments by varying the polar angle from 25° to 155°, for two azimuthal angles corresponding to the (-110) and (11-2) planes. Since the kinetic energy of the photoelectrons are well above 1 keV, we restrict ourselves to the simplest interpretation of the photodiffraction data, relying on the predominance of the forward scattering at such energies [51] and considering single diffraction only. Within this simple picture, the different CL intensity profiles mainly reflect the diffracted photoelectron angles calculated from interatomic directions within a 2D projection in the (-110) and (11-2) planes (See S3, Fig. S3.1(a,b), S3.2(a,b) and Tables S3.I-IV). The recorded Ga 3d and P 2p CL intensity profiles on the bare surface and the Se exposed one are shown in Figures S3.1(c,d) and S3.2(c,d) for the polar angle varying with respect to the [11-2] direction (i.e. within the (-110) plane) and [1-10] direction (i.e. within the (11-2) plane) respectively. For the bare surface, the measured spectra are indeed in good agreement with the predictions of the forward scattering model with some slight discrepancies, which are out of the scope of this work. In the case of the photodiffraction within the (-110) plane, the spectra before and after Se exposition are rather similar with a decrease of the P 2p intensity, which could indicate the replacement of P atoms by Se ones on the surface. Turning now to the case of the diffraction within the (11-2) plane, the Ga 3d and P 2p spectra on the bare surface are not symmetrical, contrarily to what is expected for the (1 × 1) surface reconstruction corresponding to the model used for the evaluation of the diffraction angles. This asymmetry could be ascribed to the (2 × 2) reconstruction of the deoxidized GaP surface, due to the presence of P-P bonds at the surface. After Se exposure, more symmetrical variations are recovered, in good agreement with the (1 × 1) diffraction pattern observed in this case. Once again, the intensity of the P 2p CL after Se treatment is strongly reduced, confirming the replacement of P atoms by Se ones on the surface. A further detailed description of the slight changes in the various spectra upon Se exposure is presented in Supplementary S3 but the main information comes from the superposition of the P 2p, Ga 3d and Se 3d CLs intensity profiles for the two selected azimuths (Fig. 5(a) and 5(b)). Indeed, these figures highlight the somewhat flat selenium profile without any prominent peak, which means that photoelectron emitted by Se atoms do not experience any forward scattering before leaving the surface. This is a strong evidence that Se atoms mainly lie on the top surface atomic layer. Yet, the normalization of the Se 3d profile within the (11-2) plane (Fig. 5(c)) reveals three slight peaks at 75°, 90° and 115°, which might be an indication of the presence of Se atoms in the second atomic plane but to a low extent. Nevertheless, the present selenium treatment never leads to a $Ga_xSe_y$ surface phase as evidenced in the work of Bezrvadin et al. [31] who mentioned the formation of $Ga_2Se_3$ after GaP(111) passivation under strongly higher selenium partial pressures than the ones we use.
<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /></figure>

<figure><img src="image_9.png" /><figcaption>FIG. 5. Photodiffraction intensity profiles of P 2p, Ga 3d and Se 3d CLs, extracted from a GaP(111)B substrate passivated at 700°C under $F_{Se}=3.4×10^{-6}$ Torr within the (a) (1-10) and (b) (11-2) atomic planes. Normalized photodiffraction intensity profile of Se 3d CL within the (11-2) plane(c).</figcaption></figure>

In order to reduce the sub-surface component evidenced in the Se 3d CL line shape in Fig. 4(c), we have varied both $F_{Se}$ and $T_{Se}$, considering two ranges for $F_{Se}$, 3-5×10⁻⁷T or 3-5×10⁻⁶T, and two values for $T_{Se}$, 450°C or 700°C. Figures S4(a) and S4(b) show that the Se3d/Ga3d and Se3d/P2p CL intensity ratios decrease with increasing take-off angles, confirming that selenium lies mainly at the top surface in all cases. More, the final annealing temperature $T_{Se}$ has little influence on the chemistry of the final surface whereas it is essentially dependent on the selenium flux $F_{Se}$. This is further confirmed by the Se 3d CL line shapes reported in Figures 6(a) and 6(b) for the four combinations of $F_{Se}$ and $T_{Se}$ considered. Indeed, the subsurface component is only prominent for the two
spectra recorded for the highest $F_{Se}$ range, whatever $T_{Se}$. The decomposition of the Se 3d CL for the two ranges of $F_{Se}$ at $T_{Se} = 450^{\circ}$C (Fig. S5) highlights a strong reduction of the subsurface component when decreasing $F_{Se}$. Figure S5(a) is strongly similar to Figure 3(h), further evidencing the little effect of the maximum temperature exposure on the resulting Se 3d line shape.

For the lowest selenium flux $F_{Se}$, i.e. with the minimum sub-surface component, we then investigate the resulting variations of the band bending and surface dipole for the $T_{Se} = 450$ and $700^{\circ}$C. To evaluate the band bending after Se exposure, we rely on the binding energy variation $\Delta E_{P2p}$ of the P 2p$_{3/2}$ CL bulk component with respect to the bare surface since it gives directly the variation of the surface band bending. The variation of the surface dipole is obtained through work function measurements as for the bare surface. Figure S6(a) evidences a systematic shift of about 0.2 eV to higher binding energy of the P 2p CL, indicative of a band bending reduction and hence of a partial electronic passivation. Figure S6(b) shows a decrease of about 0.1 eV of the threshold energy, indicative of an increase of 0.1 eV of the work function of the Se-exposed surface. Together with the band bending decrease, this leads to a change of the surface dipole from negative values around -0.2 eV on the bare surface to positive values around +0.1-0.15 eV on the Se-exposed surface. The positive values of $\Delta\Phi_s$ upon Se exposure is indicative of the presence of negative charges on the outer surface, while positive charges are located on the inner one. It is in line with the stronger Se electronegativity than the P one, which is a supplementary evidence of the top surface Se passivation. Table II summarizes the data and Figures 6(c) and 6(d) give a sketch of the surface electronic structure after Se exposure. Taking into account the 0.2 eV band bending reduction, the VBM energy lies around -1.8 eV with respect to the Fermi level, i.e. about 0.2 eV higher than the value previously determined by ARPES (-2.0 eV, Fig. 4). This discrepancy is assumed originating from the different methods used to determine the VBM in ARPES and in UPS.
<figure><img src="image_10.png" /><figcaption>FIG. 6. Evolution of the Se 3d CL with $F_{Se}$ for a maximum exposure temperature $T_{Se}$ of 700°C (a) or 450°C (b). Band structure schematic of Se-passivated n-type $GaP(111)_B$ substrate with (d) $T_{Se}=450°C$ and (e) $T_{Se}=700°C$ for a $F_{Se}$ flux in the $3-5 \times 10^{-7}$ Torr range.</figcaption></figure>

<table><caption>TABLE II. Comparison of the surface electronic properties between a deoxidized and Se passivated n-type $GaP(111)_B$ surface for 2 Se exposure temperatures ($T_{Se}$) and $F_{Se}$ in the $3-5 \times 10^{-7}$ Torr range. Measured P 2p binding energy (P 2p BE) and work function (Φ); calculated variations of the band bending ($Δ_{BE}$=-ΔΦ<sub>b</sub>), the work function (ΔΦ) and the surface dipole (ΔΦ<sub>s</sub>). Deduced values of the band bending (Φ<sub>b</sub>) and of the surface dipole (Φ<sub>s</sub>).</caption><thead><tr><th>Sample</th><th>P 2p BE<br>(eV)</th><th>Φ<br>(eV)</th><th>Δ<sub>BE</sub>=-ΔΦ<sub>b</sub><br>(eV)</th><th>ΔΦ<br>(eV)</th><th>ΔΦ<sub>s</sub><br>(eV)</th><th>Φ<sub>b</sub><br>(eV)</th><th>Φ<sub>s</sub><br>(eV)</th></tr></thead><tbody><tr><td>n-type GaP(111)<sub>B</sub></td><td>129.75</td><td>4.31</td><td></td><td></td><td></td><td>0.63</td><td>-0.19</td></tr><tr><td>GaP(111)-Se T<sub>Se</sub>=450°C</td><td>129.99</td><td>4.42</td><td>0.24</td><td>0.11</td><td>0.35</td><td>0.37</td><td>0.16</td></tr><tr><td>GaP(111)-Se T<sub>Se</sub>=700°C</td><td>129.94</td><td>4.43</td><td>0.19</td><td>0.12</td><td>0.31</td><td>0.43</td><td>0.12</td></tr></tbody></table>

### C. Growth of few layers TMDCs on Se-terminated $GaP(111)_B$

Recently, our group has demonstrated the large scale growth of few layers epitaxial $WSe_2$ on $GaP(111)_B$ after a proper Se-passivation of the substrate. We underscored an optimal growth temperature between 450°C and 500°C followed by a subsequent annealing step under a Se flux in order to promote the semiconducting 2H phase [52,53], while evidencing the formation of the 3R phase at 250°C [54]. In addition, we also demonstrated the large scale growth of $TaSe_2$ with a strong dependence of the growth and annealing temperature on the 1H or 1T phase formation [55]. The epitaxial growth of monolayer 1T-$TaSe_2$ is of great interests in the development of large-scale 2D Mott material, with the observation of a Mott insulating gap up to 400K [56]. In the following, we present few
results on the quasi-van der Waals epitaxial growth of HfSe₂ on a Se-terminated GaP(111)B substrate developed at 250°C and followed by a subsequent annealing step at 680°C under selenium flux. The HR-STEM image in the high HAADF mode along the (110) zone axis (Figure 7a) evidences a clear van der Waals gap at the interface. As the brightness contrast is approximatively proportional to Z², the observed contrast reveals an extreme GaP surface selenium passivation before growth. Overall, we point out a good substrate to TMDC interface. The HfSe₂ multilayer exhibits a single 1T-phase, as shown by the blue lines, with numerous changes in the grain orientation within two subsequent layers. We determine a [1120]$_{HfSe_2}$//[110]$_{GaP}$ and [1100]$_{HfSe_2}$//[112]$_{GaP}$ in-plane epitaxial relationship, and a [0001]$_{HfSe_2}$//[111]$_{GaP}$ out of plane relationship.

<figure><img src="image_11.png" /><figcaption>FIG. 7. (1 × 1) RHEED pattern and HAADF-STEM image of a HfSe₂/Se-GaP(111)B heterostructure grown with MBE (a) and the associated side view schematic created with VESTA (b).</figcaption></figure>

## IV. CONCLUSION

In summary, we have demonstrated the preparation, structural and electronic properties of the GaP(111)B surface, from its deoxidation to the Se-termination required for the TMDC growth. We have underscored the need for a combination of AH and cracked PH₃ rather than thermal deoxidation, while evidencing the key role the AH exposure time and a maximum substrate temperature of 550°C to a successful deoxidation. The resulting surface exhibits a (2 × 2) reconstruction induced by a trimer periodic arrangement of phosphorus atoms in excess on the surface and clear atomic steps with serrated edges. Under the exposition of a selenium flux, we have demonstrated a change of the GaP(111)B surface reconstruction to a (1 × 1) pattern, consistent with the suppression of the phosphorus excess. We have highlighted a similar morphology before and after the Se-treatment, while evidencing that selenium mainly reacts with the P-terminated surface without forming a GaSe layer. On the other hand, XPS and UPS measurements have demonstrated an effective partial electronic passivation of the surface through the evolution of both the band bending and the surface dipole. ARPES measurements also point out that the band structure of the Se-passivated surface originates from the GaP(111)B one without any change associated with selenium. However, the observation of residual selenium atoms in the second atomic plane has required the optimization of the process, unveiling the predominant impact of a selenium flux in the 10⁻⁷ Torr range, whereas the substrate temperature has shown little effect between 450°C and 700°C. This first realization of Se-terminated GaP(111)B substrate could pave the way to new 2D/3D heterostructure development with great interests in nanoelectronic and optoelectronic applications.

## ACKNOWLEDGMENTS
We thank Dr. Ludovic Desplanque for fruitful discussions about III-V (111)-oriented surfaces. We acknowledge the financial support by the Tunne2D (ANR-21-CE24-0030) and ADICT (ANR-22-PEEL-0011) projects, as well as the French technological network RENATECH and Region Hauts-de-France. D.R. acknowledges support from the HPC resources of IDRIS, CINES, and TGCC under Allocation No. 2024-A0160914101 made by GENCI.

**Competing financial interests:** There are no conflicts to declare.

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<figure><img src="image_12.png" /></figure>