Keywords: EPR spectroscopy, molecular nanomagnets, spin chirality, spin-electric coupling, spin triangles

First Demonstration of Magnetoelectric Coupling in a Polynuclear

Molecular Nanomagnet: Single-Crystal EPR studies of [Fe3O(O2CPh)6(py)3](ClO4)•py under static electric fields Athanassios K. Boudalis,* [a] Jérôme Robert, [a,b] and Philippe Turek [a] Abstract: Single-crystal EPR experiments show that the highly symmetric antiferromagnetic half-integer spin triangle [Fe3O(O2CPh)6(py)3](ClO4)•py (1) possesses a ST = 1/2 ground state exhibiting high g-anisotropy due to antisymmetric exchange (Dzyaloshinskii-Moriya) interactions. EPR experiments under static electric fields parallel to the triangle's plane (i.e. perpendicular to the magnetic z-axis) reveal that this ground state couples to externally applied electric fields. This magnetoelectric coupling causes an increase in the intensity of the intradoublet EPR transition and does not affect its resonance position when B0||z. The results are discussed on the basis of theoretical models correlating the spin chirality of the ground state with the magnetoelectric effect.

Electric control of magnetic materials promises to revolutionise the design of electronic devices, as it entails levels of spatial and temporal precision unattainable through magnetic fields. Sources of electric fields, such as STM tips, can easily address single molecules, or atoms, and implement switches that can turn on and off at rates of several tens of GHz. Such specifications are beyond the reach of magnetic field sources. However, this magnetoelectric control usually requires a mechanism that can couple spins to electric fields, such as the hyperfine Stark effect reported in [Tb III (pc)2] (pc -= phthalocyanine anion). [1] Interestingly, antiferromagnetic triangles of half-integer spins, or "spin triangles", are magnetic objects that lack an inversion center, and whose spin should, in principle, directly couple to external electric fields. In addition, those objects possess a rare quantum property, spin chirality, which derives from the simultaneous interplay of competing antiferromagnetic interactions and Dzyaloshinskii-Moriya interactions (DMI). Spin chirality has generated interest as it relates to diverse phenomena, such as orbital currents in Mott insulators, [2] toroidal magnetic moments in Dy III 3 triangles, [3] Berry phase in spin triangles, [4] magnetic skyrmions [5] and anomalous Hall effect in chiral magnets. [6] The scalar spin chirality in spin triangles has recently been proposed for the implementation of spin-chirality qubits, utilizing the eigenvalue of their spin chirality operator as a computational degree of freedom, instead of their spin projection. This encoding scheme, initially proposed for Molecular Nanomagnets (MNMs) [7,8] and then extended to clusters of atomic Cu [9] and Quantum Dot triads, [10] promises electric control, [7,8] practically zero decoherence from hyperfine interactions, [11] and even the possibility of quaternary logic. [12] Moreover, apart from its use in single-qubit control, the postulated magnetoelectric effect has also been proposed for long-range multi-qubit coupling by the microwave electric field inside resonant cavities. [13] In the proposed schemes, [4,7,8,13,14] an electric field E parallel to the triangle's plane (i.e. perpendicular to the molecular magnetic zaxis, Ez) is proposed to couple to the spin of the ground state, allowing its electric control.

The high availability of molecular triangular magnetic systems, such as 'basic' iron/chromium(III) carboxylates and copper(II) pyrazolates, [15,16] and their demonstrated stability upon surface deposition, [17] render them particularly attractive for a chemical approach to the problem of qubit implementation. EPR spectroscopy is a powerful tool towards confirming the feasibility of this approach; EPR studies under electric fields have recently been applied for the electric control of the spin states of Mn II in ZnO, [18] 75 As and 31 P in Ge and Si, [19][20][21] while Ferromagnetic Resonance (FMR) has been used to study the electric control of multiferroic heterostructures. [22] We had previously carried out [23,24] extended magnetic, Xband EPR and low-temperature synchrotron crystallographic studies on the highly-symmetric complex [Fe3O(O2CPh)6(py)3](ClO4)•py [25] (1, Figure 1). Herein, we report X-band CW-EPR spectroscopic studies on single crystals of 1 under externally applied static electric fields and demonstrate a magnetoelectric coupling for the first time in an exchangecoupled molecule and in MNMs in general. Single-crystal EPR spectra for 1 ( Figure 2) are in line with the powder ones: they reveal a symmetric derivative spectrum at g = 2.0 when the molecular axis (z) is parallel to the Zeeman field (B0||z), which corresponds to the g0|| value of the constituent ions (Oh high-spin Fe III ). The resonance shifts to higher fields when the angle increases, accompanied by: (i) a pronounced decrease of the spectral amplitude, (~30:1 between the 0 and 90° spectra, respectively), (ii) a broadening of the lines and (iii) the appearance of spectral features which we have previously attributed to atomic vibrations which lead to magnetic symmetry lowering, in conjunction with DMI. [24] Single-crystal spectra were collected under static electric fields (E0) using a custom-made sample holder, suitable for brittle molecular crystals of arbitrary shapes, for use on commercial instruments (see Supporting Information). The field was applied parallel to the plane of the Fe3 triangles as considered by Loss [7,8,11,13] and others [4,14] (i.e. E0z), and rotated along with the crystal on the horizontal laboratory plane (Scheme S1). Considering an electrode distance of r ~ 0.25 mm, the maximum electric field for a voltage of ΔV = 2200 V, was ΔV/r ~ 9×10 6 Vm -1 (90 kV/cm).

For B0||z, it was readily observed that the intensities of the spectra increased upon application of the electric field (Figure 3). For the quantification of the effect, fits with a Voigtian lineshape were carried out for each spectrum (Figure S1). Based on those fits, the increase in amplitude was associated with an increase of intensities of the spectra (double integrals) and a decrease of their linewidth's Gaussian component, while the Lorentzian component remained stable; the g-factor was practically unaffected, within experimental error, and equal to 2.0010±0.0005. The integrated intensities of the experimental data increase linearly above 500 V (Figure 3, inset).

At a 45° (B0,z) angle the spectra exhibited a similar effect (Figure 4). Since the line shape in this case was more complicated and non-symmetric, fits were not attempted. Instead, the intensities of the experimental curves (double integrals) were extracted and compared. In comparison, between 0 and 2200 V, the 45° spectra exhibit a slightly less pronounced relative increase of their intensity (Table 1); this increase appears linear within experimental error (Figure 4, inset). The effect was less discernible for experiments carried out at B0z (Figure 5). The main reason for this was that, due to experimental considerations (see Supporting Information), a dielectric resonator was used for studies under electric fields, whereas a high-sensitivity TE011 cavity was used for standard single-crystal and powder studies. The lower sensitivity of the dielectric resonator, combined with the smaller spectral amplitudes at B0z (~3% of those at B0||z), resulted to significantly lower S/N ratios. Double integration of the spectra indicated a +8% increase in the spectral intensity, although this figure should be considered with caution given the lower sensitivity of these experiments. The magnetic exchange within complex 1 and similar complexes are described by the general spin Hamiltonian where Jij are the isotropic (Heisenberg) exchange parameters and Gij are the pseudovectors of the antisymmetric (Dzyaloshinskii-Moriya) exchange terms and which for triangular complexes are characterized by Gz >> (Gx, Gy) ~ 0 as predicted by Moriya's symmetry rules; [26] single-ion zfs does not appreciably affect the EPR spectra of these complexes even for high single-ion D values, as we have previously shown. [24] Theoretical models have been proposed to account for the effects of static electric fields on the EPR spectra of triangles experiencing DMI. Detailed such models have been developed for triangles of Si = 1/2 spins [8,27,28] and have been briefly touched upon for triangles of Si = 3/2 spins, [START_REF] Trif | Spin-Electric Coupling in Quantum Dots and Molecular Magnets[END_REF] but none have been explicitly worked out for Si = 5/2 triangles. Loss and coworkers [8] have worked out the Stark effect on the frequencies and intensities of EPR resonances of equilateral triangles (J12 = J23 = J13 = J) for E0z, and in two Zeeman orientations B0||z and B0z. Belinsky [27,28] has worked out the influence of an E0z field on equilateral, isosceles (J12 = J23 = J  J23 = J') and scalene (J12  J23  J13  J12) triangles, and has given a detailed description of the relations between vector chirality 𝐊 ̂, the scalar chirality 𝐂 ̂, the spin configurations and magnetizations in all those cases. In particular, for an isosceles triangle in B0||z, the four low-lying states are divided into two doublets, with the ground doublet being characterized by an intermediate spin S12 = 0 (for |J| < |J'|) or S12 = 1 (for |J| > |J'|), and by a right-handed vector chirality (κ = +1) for Gz < 0 or left-handed chirality (κ = -1) for Gz > 0. Each Zeeman sublevel is characterized by scalar chirality χ = 2κMs = ±1 (Figure S2). Therefore, the classical EPR intradoublet transition conserves the vector chirality and the intermediate spin, and reverses the scalar chirality (Δχ = ±2) and spin projection (ΔMs = ±1). For Gz = 0 the system is achiral (χ = κ = 0), while for an equilateral system (J = J') the states possess no S12 character. [28] Focusing on the ground |χ,Ms doublet of an equilateral triangle, Loss and coworkers [8] have calculated the effect of an electric field E0z on the intensity of the intradoublet |-½,+1 ↔ |+½,-1 EPR transition as proportional to:

|-½,+1|Sx|+½,-1| 2 = |-ε 2 /D +1 D -1 | 2 (2) where ε = d•E and D χ = {ε 2 + [Δ + χ(ε 2 +Δ 2 ) 1/2 ] 2 } 1/2 .
In these expressions d is the dipolar coupling constant and Δ is the energy separation between the two doublets, which is considered to be solely due to DMI and not to any magnetic asymmetry. Equation (2) implies zero transition probability at E0  = 0, which indeed is the case for an equilateral triangle; [START_REF] Rakitin | [END_REF]31] in that sense, the role of a E0z field is to break the equilateral symmetry by modifying the Jij parameters. For lower magnetic symmetries (ΔJ = J -J'  0) this is an allowed EPR transition even for E0  = 0, and Belinsky shows [28] that the increase of ΔJ would increase its intensity. Ab initio calculations by Islam et al. [14] on an equilateral Cu II 3 polyoxovanadate have indicated that a field E0z of ca. 10 7 Vm -1 , i.e. comparable to ours, would increase ΔJ by ca. 2×10 -3 cm -1 ; this, in turn, would increase the intradoublet transition intensity, according to Belinsky. However, Belinsky provides no analytical expressions for the EPR transition probabilities as a function of the applied electric field in the general case of an asymmetric triangle. As for the resonance position, both Loss [8] and Belinksy [28] predict that the intradoublet B0||z transition will not shift in energy under the influence of an electric field E0z.

From a qualitative point of view, our observations confirm the theoretical predictions of Belinsky and of Loss. There is indeed a magnetoelectric coupling as predicted from symmetry arguments. Also as predicted, the intradoublet transition increases in intensity with the application of an electric field E0z, and its resonance position remains unshifted when B0||z. This increase in intensity is in line with an increase in magnetic deformation, though other mechanisms may also be operative. It should also be noted that, theoretically, such a decrease in

, 1 , 1 ˆˆˆ2 2 ( ) (1
) ij i j ij i j i j i j HJ        S S

G S S

magnetic symmetry would also be reflected on the position of the transition at B0z, as the effective g is determined by the relation between ΔJ and Gz. [START_REF] Rakitin | [END_REF]31] However, the lower sensitivity at that orientation does not allow us to make that determination. Finally, our observations regarding the line shapes upon application of electric fields, i.e. narrowing and increase in peakto-peak amplitude, are not predicted by either model at the considered level of precision.

The EPR signal response to an external electric field should, in principle, allow the direct determination of the electric dipole coupling constant, d, for our complex. However, quantitative analyses are not currently possible, as the case of a Si = 5/2 triangle has not been worked out.

In addition, the electric field effect would need to be explicitly considered in conjunction with the dielectric constant of the crystal. While the above mentioned models refer to effective electric fields experienced by isolated molecules, molecules in real crystals will tend to experience a significantly different electric field than calculated by ΔV/r. We therefore need to consider both the depolarizing field that will reduce the macroscopic field within the crystal, as well as the local Lorentz field experienced by the molecular sites in the crystal. Since, however, the material's dielectric constant cannot be assumed, it must be determined for the particular experimental conditions. Indeed, it is known that the dielectric constant of a given material depends on the temperature and the electric field frequency, and it was recently shown that in the case of two molecular ferromagnets, it can also heavily depend on the applied magnetic field. [32] Therefore, to safely quantify the applied field on the molecular sites, precise such determinations are necessary for 1.

The reported magnetoelectric coupling is a critical prerequisite for the electric control of spin qubits. In the present context it affords preliminary confirmation to the prediction that oscillating electric fields can induce scalar chirality inversions in spin-chiral systems, thus encoding spin qubits with long decoherence times. In the broader context of spin-electric coupling in MNMs, it demonstrates that direct magnetoelectric coupling is possible based on symmetry considerations alone, without the need of additional mechanisms, such as hyperfine interactions, thus broadening the scope and applicability of the effect.

In conclusion, we have confirmed the magnetoelectric coupling predicted for spin triangles under electric fields. This is, actually, the first such experimental confirmation in molecular exchange-coupled polynuclear systems, for the purposes of which we developed a custom sample holder. We are currently working toward increasing the maximum electric field, to reach beyond the linear region of the signal increase, and toward experiments at other orientations of the electric field with respect to the molecular z-axis (e.g. Ε0||z). We are also working towards developing theoretical models for the treatment of Si = 5/2 triangles and determining the dielectric constant of 1 to accurately determine the electric fields at the molecule sites inside the crystal. At the same time, we are planning similar studies on Cu II 3 and Cr III 3 triangles, for which the theoretical framework is better developed.

Figure 1 .

 1 Figure 1. POV-Ray plot of the cation of 1. The green arrow passing through the central oxide indicates the crystallographically imposed C3 axis, running parallel to the crystallographic c-axis.

  K. Boudalis, Prof. Philippe Turek Institut de Chimie de Strasbourg (UMR 7177, CNRS-Unistra), Université de Strasbourg, 4 rue Blaise Pascal, CS 90032, F-67081 Strasbourg, France. E-mail: bountalis@unistra.fr [b] Dr. Jérôme Robert Sorbonne Université, CNRS, Laboratoire Jean Perrin, LJP, F-75005 Paris, France Supporting information for this article is available on the WWW under http://dx.doi.org/10.1002/chem.2018xxxxx.

Figure 2 .

 2 Figure 2. X-band EPR single-crystal and powder data of 1 at 4.2 K. The asterisk at ~3300 G indicates a cavity impurity. Experimental conditions: f = 9.31 GHz, mod. ampl. = 10 Gpp. Single-crystal spectra (0-90°): Power = 2.5 mW. Single-crystal spectrum at 90° (red line): Power = 10 mW. Powder spectrum (black line): Power = 0.02 mW.

Figure 3 .

 3 Figure 3. X-band EPR single-crystal data of 1 at B0||z under varying electric fields. Experimental conditions: f = 9.80 GHz, mod. ampl. = 10 Gpp, power: 2 mW. The inset shows the electric field dependence of the integrated intensities of the experimental spectra (Imin = 1).

Figure 4 .

 4 Figure 4. X-band EPR single crystal data of 1 at (B0,z) = π/4 under varying electric fields. Experimental conditions: f = 9.80 GHz, mod. ampl. = 10 Gpp, power: 4 mW. The inset shows the electric field dependence of the integrated intensities of the experimental spectra (Imin = 1).

Figure 5 .

 5 Figure 5. X-band EPR single crystal data of 1 at B0z inside and outside an electric field. Experimental conditions: f = 9.80 GHz, mod. ampl. = 15 Gpp, power: 6.2 mW.

  33

Table 1 .

 1 Relative changes of integrated intensities between 0 and 2200 V.

		0°	45°	90°
	Intensity	+13%	+12%	+8%

Acknowledgements

We thank Prof. Daniel Loss for helpful discussions on theoretical aspects and experimental design, Dr. Yiannis Sanakis for discussions on the EPR spectroscopy of spin triangles and the CNRS for the temporary assignment of JR at the Institut de Chimie de Strasbourg for the duration of this work. This project has received funding from the European Union's Horizon 2020 research and innovation programme under the Marie Sklodowska-Curie grant agreement No 746060 (project "CHIRALQUBIT").

Entry for the Table of Contents

COMMUNICATION

Clear! An electric field applied on the antiferromagnetic spin-chiral complex [Fe3O(O2CPh)6(py)3](ClO4)•py couples to the spin of its ground state and modifies its Electron Paramagnetic Resonance (EPR) spectrum. This magnetoelectric coupling can provide electric control of the spin states of Molecular Nanomagnets.


T Demonchaux

K K Sossoe

M M Dzagli

J P Nys

M Berthe

D Troadec

A Addad

M Veillerot

G Patriarche

H J Von Bardeleben

M Schnedler

C Coinon

I Lefebvre

M A Mohou

D Stiévenard

J F Lampin

Ph Ebert

X Wallart

B Grandidier

Chemical nature of the anion antisite in dilute phosphide GaAs1-xPx alloy grown at low temperature

While non-stoichiometric binary III-V compounds are known to contain group V antisites, the growth of ternary alloys consisting of two group V elements might give additional degrees of freedom in the chemical nature of these antisites. Using cross-sectional scanning tunneling microscopy (STM), we investigate low-temperature grown dilute GaAs1-xPx alloys.

High concentrations of negatively charged point defects are found.

Combined with transmission electron microscopy and pump-probe transient reflectivity, this study shows that the defects have a behavior similar to the group V antisites. Further analyses with x-ray diffraction point to the preferential incorporation of arsenic antisites, consistent with ab-initio calculations, that yield a formation energy 0.83 eV lower than for phosphorus antisites. Although the negative charge carried by the arsenic antisites in the STM images is shown to be induced by the proximity of the STM tip, the arsenic antisites are not randomly distributed in the alloy, providing insight into the evolution of their charge state during the growth.

I. INTRODUCTION

Molecular beam epitaxy (MBE) is widely used for the epitaxial growth of III-V compounds. This technique has been essential to obtain ternary or quaternary alloys with intermediate lattice parameters and band gaps between those provided by binary compounds. Tuning these properties is commonly achieved by changing the fraction of the alloyed atoms in the material based on rules that proved their worth. For ternary alloys with two different group III elements, the composition of the semiconductor is thermodynamically controlled and is predominantly governed by the input flux ratio of both elements.

As to ternary alloys containing two different group V elements, the composition is highly nonlinear with the fraction of the group V incident flux ratio and it is usually tailored with the temperature of the substrate. Deviations from these rules of thumb exist [1], but all these deviations have so far been understood and included in thermodynamic analyses of the growth to successfully predict the composition of ternary alloys depending on the growth parameters.

While the MBE growth of III-V semiconductor usually yields stoichiometric material, since group V elements stick to the surface only if group III atoms are present, it has been made possible to grow nonstoichiometric binary compounds with high crystal quality by lowering the substrate temperature [2]. In this case, group V elements are also incorporated on cation lattice sites, giving rise to group V antisites [3,4,5]. Such a low temperature growth regime has been much less investigated for ternary compounds [6,7,8, 9 ] and it is an open question whether the excess of anion atoms incorporated on the cation sublattice follows the same fraction of group V elements incorporated on the anion sublattice. Because antisites are known to introduce defect states in the band gap of materials that trap free charge carriers and pin the Fermi level, controlling their chemical nature and thus the energy of the defect states would provide an additional degree of freedom in the generation of electromagnetic waves with ternary III-V compounds in the THz frequency range [10,11,[START_REF] Currie | Photodetectors: Materials, Devices and Applications[END_REF][START_REF] Bacon | [END_REF]]. An appropriate choice of the ratio of group V elements could also help in tuning the strain of the epilayer due to the incorporation of excess group V elements, allowing for the growth of layers with a wider range of critical thicknesses [14,15].

Here, we address this fundamental question and study the chemical nature of the antisite in a nonstoichiometric ternary compound semiconductor, that consists of a dilute phosphide GaAs1-xPx alloy. This alloy is a prototype system, since the mismatch in bond stiffness, lattice constant and orbital energy between GaAs and GaP is rather small. The incorporation of excess group V elements is achieved by growing a dilute phosphide GaAs1-x Px alloy with molecular beam epitaxy at low temperatures ranging between 240°C and 325°C. From the direct visualization of the native defects with cross-sectional scanning tunneling microscopy (STM), we identify the majority of the point defects as antisites, despite characteristic features that depart from the ones encountered in previous STM studies [3,16,17,18]. In particular, the antisites are found to be negatively charged. Additional grounds for the identification of the point defects as antisites are provided by pump-probe transient reflectivity and the analyses of annealed samples with STM and transmission electron microscopy (TEM).

Characterisation of the strain in as-grown and annealed samples by high-resolution x-ray diffraction reveals the formation of a dilute phosphide GaAs1-xPx alloy, that contains a high concentration of arsenic antisites (AsGa). Ab-initio calculations of the formation energy of arsenic antisites and phosphorus antisites (PGa) show a lower formation energy for AsGa, supporting the incorporation of excess arsenic into the epilayer instead of phosphorus. Finally, the charge state of the antisites was investigated in the low-temperature grown (LTG) layers. We show the key role of the tip induced band bending and the temperature to account for the observation of negative charge states close to the cleaved (110) surface. Based on electron-paramagnetic-resonance (EPR) analyses of the samples and the study of the antisite distribution in the STM images, we are also able to trace back the charge state of the AsGa antisites during their incorporation into the LTG layers.

II. EXPERIMENTAL AND THEORETICAL DETAILS

1.5 µm-thick LTG layers were grown on n-type GaAs (001) substrates (3x10 18 cm -3 ) by molecular beam epitaxy (MBE). They were separated from the substrates by a 50 nm-thick GaAs buffer layer and a 100 nm-thick Ga51In49P potential barrier used as a marker to locate the LTG layer in the STM experiment. During the formation of the LTG GaAs layer, a small amount of P atoms was incorporated to obtain a dilute phosphide alloy. The concentration of P in all the samples that were prepared did not exceed 2.3 x10 20 cm -3 , thus corresponding to a ratio between P atoms and As atoms of ~1% or less. For all the samples, such a concentration was analyzed with SIMS to provide elemental depth profiles along the [001] growth direction (see Fig. 1). Different growth temperatures (Tg) between 325°C and 240°C were used to change the excess of group V elements during the formation of the LTG layers. For all samples, the growth temperature of the LTG layers was measured by band-edge absorption spectroscopy. After growth, the samples were cut in several pieces. Some of them were annealed. The rapid thermal annealing consisted of a temperature ramp set at 25°C/s, a stabilization of the temperature (Ta) at 580°C or 620°C for 40 s and a cooling to room temperature in 80 s. All these steps were performed under a flow rate of 250 sccm of diazene (N2H2).

For investigations by cross-sectional transmission electron microscopy (TEM), thin lamellae were cut from the samples by focused ion beam and deposited on TEM grids. A typical cross-sectional high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) shows the structure of interest (inset of Fig. 1). The samples were also characterized by cross-sectional STM on clean cleavage surface. For optimal cleavage in ultrahigh vacuum (UHV), the samples were first thinned back to a thickness of 150 µm and then loaded into the UHV chamber. Directly after cleavage, the freshly exposed (110) surface of the samples was immediately examined with tungsten or platinum tips in a scanning tunnelling microscope working at a temperature of 77 K. Tunneling spectroscopic measurements were performed at constant tip-sample separation to be compared with numerical simulations of the tunneling current. These simulations were based on the Semitip code [19], where the tip-induced band bending was calculated with a finite-element method assuming a hyperbolic shaped probe tip. Due to the small concentration of phosphorus in the LTG layer, the semiconductor was considered as pure GaAs, treated in an effective-mass approximation and tunnel currents were computed using the Bardeen method. The parameters in the computation were the tip radius-of-curvature, the tipsample distance, the contact potential (difference between the work functions of the tip and the sample) and the doping level in the semiconductor. We note that all the STM and tunneling spectroscopic measurements reported in this manuscript were performed in the middle part of the 1.5 µm-thick LTG layer.

In order to measure the antisite concentration, the LTG layers were also studied by highresolution x-ray diffraction. The incorporation of excess group V atoms into the crystal results in lattice strain [4,20]. Therefore, the lattice parameters relative to the substrate were obtained from rocking curve measurements, where the diffraction peak from the LTG layers was compared to the Bragg maximum from the GaAs substrate to generate the splitting parameters Δθ. The relative lattice parameter change Δa/a induced by the antisites was deduced from Δθ taking into account the contribution due to the alloy composition. Such a contribution was obtained from the knowledge of the alloy composition on the basis of the SIMS profiles.

The EPR measurements were performed with a Bruker X-band spectrometer with standard field modulation at 100kHz. The spectra were taken at T=4 K and T=16 K which are the optimal temperatures for the observation of AsGa antisites. Measurements were made with detached LTG layers, which were transferred on a suprasil quartz plate in order to avoid the perturbation of the cavity quality factor by a ntype GaAs substrate. We equally studied layers deposited on a semi-insulating GaAs substrate. The AsGa + spectrum was simulated with the 'Easyspin' program and the Spin Hamiltonian parameters obtained are typical for EL2 like antisites.

Theoretical computations of the electronic structure were done within the density functional theory (DFT). Calculations were performed using the projector augmented wave (PAW) method [21,22], as implemented in the VASP code [23,24]. The Ga3d electrons were treated as core electrons. For the bulk, the optimized cell parameter was obtained to be 5.606 Å, showing a deviation of 0.84% from the experimental value. In order to obtain the formation energies of the As and P antisites, 3x3x3 supercells were used and consisted of 216 atoms. The cells were first relaxed while their size was kept fixed. The relaxation was computed with the Perdew-Burke-Ernzerhof (PBE) functional within the generalized gradient approximations (GGA) for the exchange and correlation energy [25]. A Monkhorst-Pack k-point mesh of (3 3 3) was employed. The electron wave function was expanded in plane waves up to a cutoff energy of 282.8 eV. The atomic positions were fully relaxed. The positions were optimized until the Hellmann-Feynman forces were less than 0.01 eV/Å. Then, as the computations of the formation energies of charged point defects requires a good description of the band gap energy, the fixed-structure hybrid correction (FSHC) was added to the PBE energy while applying the range-separated hybrid Heyd-Scuseria-Ernzerhof (HSE06) exchange-correlation functional to the PBE relaxed structure [26,27,28]. The HSE06 band gap of the bulk structure was computed to be 1.65eV. To evaluate the formation energies of charged antisites, we made use of the method developed by Wu et al. [29]. This method aims at eliminating the difficulties associated with the long-range Coulomb interaction of charged defects. Considering that the cell remains electrically neutral, the defects become charged by trading electrons with bands: a positively charge defect comes with an electron in the conduction bands; conversely, a negatively charged defect comes with a hole in the valence band. Thus computations for charged defects were achieved with the indication of fulfilment of bands and the formation energies were calculated based on these rules.

Calculations of the electron density were performed by modelling the GaAs (110) surface in a supercell slab geometry consisting of 4 layers of GaAs in a 5×6 surface unit cell, including 240 Ga or As atoms. A vacuum region of 8.8 Å was used for all the calculations. In such a cell, the separation between a defect and its periodic image is sufficiently large to ensure negligible interaction. The bottom layer of the slab was passivated with pseudo-hydrogen (30 atoms H1.25:1s 1.25 and 30 atoms H0.75:1s 0.75 ) to saturate the dangling bonds [30]. The atomic positions were fully relaxed in all but the bottom layer. The electron density of the surface defects was computed for slab cells at the PBE level. In order to roughly mimic the STM experiments, the density of the highest occupied levels was integrated from the two topmost layers.

III. ATOMIC SCALE IDENTIFICATION OF THE POINT DEFECTS

Figure 2 shows a typical cross-sectional STM image of the exposed (110) face of the LTG layer grown at 240°C. Numerous bright features surrounded by a dark ring-like depression that is superimposed to the atomic lattice appear in this high resolution image. The presence of all these small features spread out on the surface is consistent with the high density of bulk point defects expected for a GaAs layer grown at low temperature. However, at first sight, most of the defects do not show the common characteristics of the subsurface AsGa antisite, that consists of a central bright core and two weaker satellite peaks along the and directions [3,16,17,18].

In took into account the presence of a band of Ga dangling bond surface states, the so-called C3 empty pzlike dangling bond states, positioned 0.25 eV above the conduction band, that is found to limit the contribution of the tip-induced accumulation layer in the apparent band gap region, consistent with previous work [19].

Assuming in first approximation a similar band structure for the dilute GaAs1-xPx alloy and the pure GaAs as well as a low free carrier density in the LTG alloy (1x10 16 cm -3 ), we were able to reproduce the increase of the tunneling current above +0.7 V and below -1.5 V with a Fermi level position 0.5 eV below EC at the GaAs1-xPx cleavage surface. This Fermi energy is consistent with a lack of pinning in the bulk undoped LTG layer and a highly n-doped GaAs substrate [31], spilling some electrons into the undoped GaAs1-xPx alloy. Despite such an electron spilling, the free carrier concentration is still low and leads to a large tip-induced band bending (TIBB). At negative sample voltages this TIBB drags the conduction band edge below the Fermi level. Hence, the bottom of the conduction band is populated with electrons and these electrons can tunnel from this accumulation zone first and only at larger negative voltages electrons tunnel from filled valence band states into the tip.

This gives rise to the point of inflexion at -1.5 V (see red arrow in Fig. 2

(b)).

This interpretation is supported directly by the corrugation observed in atomically resolved STM images measured at small magnitudes of negative voltages (see e.g. inset in Fig. 2(b)). The STM image exhibits atomic rows along the [001] direction, even though the atomic separation is smaller than the ones along the direction. This orientation of the corrugation morphology occurs only for the C3 empty dangling bond state localized at the surface Ga atoms [32,33]. All other states lead to atomic rows along the direction. Hence, the tip probes at small magnitudes of negative voltage the Ga-derived C3 dangling bond states resonant with the conduction band. This finding confirms the existence of a large TIBB that also shifts the valence band edge towards more negative voltages and leads to the large apparent band gap observed in the spectrum of the LTG layer. Moreover, it highlights the importance of the Ga dangling bond states in the tunneling process at negative sample voltages, when the tip scans a well-ordered area of the LTG layer.

In the vicinity of the point defect, the contrast changes, however. It becomes darker with a We note that all these defects are also visible in the LTG layer grown at 325°C [Fig. 4(b)], albeit to a smaller concentration of 2.1 ±1.0 x 10 19 cm -3 , in agreement with the reduction of excess group V atoms as the temperature increases [37].

In order to further check whether the point defects seen in the STM images behave similarly to the antisites, some of the samples were annealed prior to the cleavage. The LTG initially grown at 240°C and annealed at 580°C for 40s exhibits no antisite anymore on its (110) cleavage surface [Fig. 4(c)].

Only dark depressions corresponding to single vacancies or clusters of vacancies are visible, along with the formation of numerous terraces. Their presence is consistent with the particular difficulty to cleave this type of sample [39]. The cross-sectional TEM analysis of the sample reveals many dark features with a mean diameter of 3.6 nm [Fig. 5]. They correspond to precipitates of group-V elements [2,38,40,41], the biggest ones giving rise to a Moiré fringe contrast (inset of Fig. 5), that is caused by a different crystal structure with respect to the LTG layer [42,43].

The existence of precipitates indicates the condensation of the point defects upon annealing. This process is expected for antisites [44]. Therefore, the disappearance of the point defects and the formation of precipitates support the identification of the point defects as antisites, although their STM contrast is slightly different from the one encountered in previous works [3,16,17,18]. We also note that, at the growth temperature of 325°C, we were not able to find precipitates in the TEM images after annealing.

Instead, the STM image of Fig. 4(d) show a lot of point defects. Despite the disappearance of the dark halo around the protrusions at a sample voltage of -2.5V, the defects still have the same characteristic satellites, allowing to categorize them as type A, B and C defects. All the defects consist of a bigger core and when both satellites are present, the type A defect has the typical features of a subsurface antisite, confirming our conclusion.

An additional characteristic feature of the antisite is the short lifetime of the photoexcited carriers when III-V LTG materials contain a high concentration of antisites [2,45]. Pump-probe reflectivity measurements were thus performed for three LTG layers grown at temperatures of 240°C, 260°C and 325°C respectively. As shown in Figure 6, a dip is seen at the lowest growth temperature of 240°C. It is caused by absorption processes due to the very high concentration of AsGa antisites [46] and results in a very short lifetime, below the ps. In contrast, at higher growth temperatures, the spectra exhibit an abrupt rise followed by a decrease of the signal. This decrease consists of a fast component at small pump-probe delay that is caused by the relaxation and cooling of the carriers in the bands. In the case of the sample grown at 325°C, it is followed by a second component that can be fitted with an exponential decay, yielding lifetimes in the range of the ps. This increase of the lifetime with the growth temperature, and hence the decrease of the antisite concentration, is in agreement with the literature [47,48,49,50]. Due to the similarity of the lifetime found in the LTG layers with the ones given in the literature for LTG binary compounds, we believe that the majority of the point defects in the bulk have the physical characteristics of the antisites. But, at the surface, they appear in three different configurations A, B, C, due to the instability of As atoms in their vicinity.

IV. CHEMICAL NATURE OF THE ANTISITE

In order to get insight into the chemical nature of the antisites, we performed XRD measurements, since this technique is able to detect the lattice expansion in a LTG layer with a significant excess of arsenic or phosphorus [4,51,52]. While the X-ray rocking curve for the LTG layer grown at the highest temperature of 325°C and a phosphorus concentration of 0.7 x10 20 cm -3 yields a single peak in the spectrum (i) of Fig. 7, it shows two peaks for the LTG layer grown with a slightly higher P concentration and the lowest temperature of 240°C (spectrum (ii) of Fig. 7). Based on the comparison of both spectra, the one centred at zero corresponds to the GaAs substrate. The other one centered at -93" reveals a strained layer that is attributed to the important excess of group-V elements in the LTG layer. As the concentration of phosphorus increases to 2.3 x10 20 P.cm -3 (spectrum (iii) of Fig. 7), the peak shifts to -30", indicating a reduction of the strain in the epilayer.

This decrease could be caused by the existence of a tensile strain in the LTG layer, when the concentration of phosphorus atoms at anions sites increases [53]. But the shift could also arise from the incorporation of PGa antisites in the epilayer that would increase if their incorporation is preferred with respect to that of AsGa antisites. Indeed, a PGa defect should induce a smaller dilated lattice distortion, since the tetrahedral covalent bond radius (r = 1.06 Å) of P is 12% smaller than that (r = 1.19 Å) of As [START_REF] Phillips | Bonds and Bands in Semiconductors[END_REF]. DFT calculations of the bond lengths in a GaAs supercell containing either a PGa antisite or an AsGa antisite confirms the change in bond radius mentioned above: a PGa-As bond has a length of dP = 2.49 Å, that is smaller than an AsGa-As bond, dAs = 2.57 Å. Such a bond length could lead to a smaller lattice expansion than the one obtained with the As antisites. In order to determine which hypothesis is valid, we analysed a sample with a growth temperature of 260°C and a P concentration of 2.2 x10 20 cm -3 . As seen in the spectrum (iv) of Fig. 7, the peak shifts to positive angles, corresponding to a lattice contraction in the epilayer. Since the decrease of the antisite concentration alone is not expected to give a peak at positive angle, such a shift is consistent with the growth of a GaAs1-xPx alloy in tensile strain on the GaAs substrate. Therefore, this result favors the incorporation of As antisites rather than P antisites in the LTG-layer.

Upon annealing, all the peaks are found at positive angles for the samples with P concentrations higher than 1.0 x10 20 cm -3 . According to the STM and TEM images shown above, the excess of group-V elements aggregate to form precipitates. As this process takes along with a release of the strain, leading to the disappearance of the lattice mismatch in pure and annealed LTG-GaAs layers [38,[START_REF] Kaminska | [END_REF]56], the observation of a shoulder is consistent again with the growth of a GaAs1-xPx LTG-layer. Assuming that the majority of the P atoms are thus located on the anion sites of the lattice, the position of the peak at positive angles should provide a good indication of the concentration of P atoms in the layers. As this contraction takes place along the growth direction only, the relative contraction (Δa/a)[001], that is measured from the peak position in the XRD measurement, is related to the relative contraction of the relaxed lattice (Δa/a)r by:

(1)

with ν = 0.31 the Poisson coefficient for GaAs [57].

Based on Vegard's law, we deduce the ratio P composition from the relative expansion of the alloy with respect to the GaAs substrate:

(2)

For the LTG layers grown at 240°C and 260°C with the highest concentration of P atoms, the XRD measurements yield a phosphorus concentration of 2.0 x 10 20 cm -3 for both samples. This concentration is in good agreement with the concentrations measured with SIMS and therefore indicates that the majority of the P atoms incorporated during the growth occupy the anion sites of the lattice.

Knowing the P concentration, we can relate the lattice expansion measured prior to the annealing to the concentration of AsGa as follows:

(

where the subscripts BA and AA stand for before annealing and after annealing, respectively. On the basis of Ref. 37, we can finally deduce the concentration of the As antisites in the LTG layers:

(4)

We find an [AsGa] concentration of 5.3 x 10 19 cm -3 for both samples grown at 240°C with two different P concentrations. The concentration is of the same order of magnitude than the concentration obtained from the count of the AsGa antisites at the cleavage surface of the same samples. This result indicates that the As atoms are preferentially incorporated as antisites rather than the P atoms.

In order to rationalize this experimental fact, a theoretical study of the formation energy of neutral, positively and negatively charged AsGa and PGa antisites was performed. Based on the method described above, the comparison of the formation energy between PGa and AsGa was achieved by computing the chemical potentials of P and As atoms in respectively the phosphorus Cmca and the hexagonal As structure limits at the HSE06 level. As shown in Figure 8, the neutral AsGa antisite is found to be more stable than the neutral PGa antisite by 0.83 eV in GaAs bulk. As the Fermi level shifts to the valence band or the conduction band, the charge state of the antisites changes. This change occurs almost at the same energy for both AsGa and PGa antisites, leading to a generally lower formation energy for AsGa antisites with respect to PGa antisites, whatever the charge state is. Therefore, the theoretical calculation supports the incorporation of AsGa antisites in a LTG dilute phosphide GaAs1-xPx alloy, consistent with the conclusion drawn from the XRD analysis.

A further hint for the preferential incorporation of As atoms as antisites is obtained by calculating the electron density in the vicinity of an AsGa or PGa antisite located in the second layer (first subsurface atomic plane) of a GaAs (110) surface. As shown in Figure 9, the electron density calculated for the neutral AsGa and PGa antisites is consistent with previous theoretical calculations [34,35,36], revealing a high electron density at the position of the antisite and along the and directions. These areas correspond respectively to the core and the satellites of the antisites seen in the STM images.

Interestingly, the PGa antisites show a reduced distribution of the electron density at the position of the satellites, no matter what the charge state is. Such a result suggests that the contrast of the satellites for the PGa antisites should be weak in the experimental STM images, opposite to the contrast observed for the antisites seen in Fig. 2, 3 and 4.

V. CHARGE STATES OF THE ANTISITE

So far, most of the STM images have revealed a depression around the AsGa antisites, meaning that the antisites carry a negative charge opposite to the neutral or positively charged state found in previous cross-sectional STM characterizations of pure GaAs LTG layers [3,16,17,18]. Although the existence of a negative charge state has been predicted [58], it has never been demonstrated experimentally and its observation casts doubt about the exact nature of the defects seen in the STM images. Therefore, the EPR technique was applied to the sample grown at 240°C with a P concentration of 2.3 x10 20 cm -3 to determine the charge state of the AsGa antisites in the bulk of the LTG GaAs1-xPx layer. Figure 10 compares the EPR spectrum measured for this sample with the spectrum of a semiinsulating GaAs substrate that is known to host a low concentration of AsGa antisites. The spectrum measured in the LTG dilute phosphide GaAs1-xPx alloy is almost flat in clear contrast with the four lines typically detected in semi-insulating GaAs. These four lines are fitted using typical values for the Lande factor g, the hyperfine constant A and the linewidth LW of each component of the quadruplet for the antisites [59] and are therefore the signature of singly positively charged AsGa antisites in pure GaAs LTG layer [60]. We used equally different optical excitations (890nm to 1200nm) in order to increase the concentration of positively charged antisites, which are the paramagnetic species. But we did not observe any photoinduced effect on the LTG layer. As our sample contains a concentration of AsGa antisites that is high enough to give signal, the absence of the peaks indicates that the AsGa antisites are neutral. This result is consistent with a Fermi level positioned in the upper part of the band gap, as already deduced by tunneling spectroscopy. Therefore, we attribute the observation of the negatively charged AsGa antisites in STM images to the proximity of the surface.

Based on the tunneling spectroscopy and the simulation of the tunneling current discussed above, it has been concluded that a strong TIBB exists when the tip scans the LTG layer. At negative sample voltages, the band bending causes the formation of an electron accumulation layer at the surface, that is directly visible through the observation of conduction band states (i.e. Ga-derived dangling bonds) in the high resolution STM images of the occupied states. The strong downward TIBB at negative voltages also shifts the AsGa antisites' charge transition level (0/-) close to the conduction band edge below the Fermi level. As a result, the proximity of the polarized STM tip changes the charge state of the AsGa antisite at the surface: a neutral antisite becoming negatively charged. This is compatible with the observed charge state in Fig. 2a.

However, at present it remains unclear, why a clear signature of a downward band bending is not always observed (see, e.g. Fig. 4d) and was not observed in prior studies of LT-GaAs by STM. As we discuss in the following, this is primarily an effect of using large negative set voltages, which shifts the dominating contribution of the tunnel current from the electron accumulation zone in the conduction band toward the filled valence band states. In order to illustrate this, the band bending was calculated and the respective band diagrams were derived at temperatures of 77 K and 300 K for different voltages (Fig. 11). The band diagrams show that the number of states in the electron accumulation zone in the conduction band changes only weakly with the magnitude of applied sample voltage. In contrast, the situation is fundamentally different for the valence band states contributing to the tunnel current: the number of valence band states increases drastically from a sample voltage of -1.6 V to -2.1 V, whatever the temperature is. Hence, with increasing magnitude of negative voltages, the weight of the valence band states in the tunnel current dominates and the depression zone around negative charge centers weakens as it is arising from spatial changes of the tunneling out of the electron accumulation zone in the conduction band. Ultimately, no signature of a charge will be detectable. This explains the absence of a depression in the STM image of Fig. 4d, that was acquired at -2.5 V. It highlights the key role of the sample voltage to slightly change the contrast of the defects. These findings allows to reconcile our observations with the STM image of Ref. 18 for a pure GaAs LTG layer grown on a n-type GaAs substrate and imaged at 77 K, where no depression was observed around the antisites.

It is interesting to note that the use of a n-type substrate makes the Fermi level unpinned in the bulk LTG layer. As the sample temperature gets lower, the Fermi level shifts to the conduction band edge. This shift takes along with a band gap broadening, leading to a reduction of the valence band states contributing to the tunnelling current for a small sample voltage [Fig. 11(a)]. This effect thus favors the observation of a depression at low temperatures for LTG layer grown on a n-type substrate. It also explains why the negatively charged antisites have never been observed for p-type LTG layers [3,16], since the Fermi level is pinned midgap in these layers [31], precluding the formation of an electron accumulation layer.

The negative charge carried by the antisite results from the shift of the transition level (0/-) below the Fermi level due to TIBB. It implies a transition level (0/-) close to the conduction band edge that is consistent with the theoretical predictions described above and the formation of a tip induced potential well with limited spatial extent [61], that decouples the antisite state from the continuum of states in the conduction band. The negative charge is also seen for AsGa antisite defects, labeled D, E and F in Fig. 2.

Based on symmetry considerations with respect to the cation sublattice, they are attributed to antisites in the 3 rd , 4 th and 5 th atomic planes. Hence, the potential well extends a few atomic planes below the surface. Deeper, in the bulk, the transition level (0/-) stays above the Fermi level and all the bulk antisites are neutral, in agreement with the EPR measurements.

Remarkably, the antisites form complex with one or two vacancies at the surface, as shown in other free As atoms and desorbing as molecules, it could be further released during the next scans to replace a vacancy, giving rise to the switching behaviour observed in Fig. 3. The negatively charged state of the antisite would therefore explain the formation of complex defects at the surface, that was not reported in previous studies of the AsGa with STM.

Finally, we notice the absence of close pairs of antisites in Fig. 2. This suggests a repulsive short range interaction. As the antisites are not seen to diffuse after the cleavage, this interaction must have been occurring during incorporation into the LGT layer, already. A quantification of this repulsive interaction is obtained by collecting the position of the antisites in STM images and calculating the distances r between all possible pairs of antisites. This operation leads to the measured probability distribution of pair distances. Dividing this quantity by the pair distances for non-interacting randomly distributed antisites yields the two-dimensional pair correlation c(r) [64], shown as inset in Fig. 12. One can recognize a region in the center around 8 nm in diameter, where the contrast indicates a low pair probability. This indicates a repulsive interaction. The pair interaction energy can be approximated in the low density limit by the mean force potential W(r) = -kT ln (c(r)). The radially integrated pair interaction energy is plotted in Fig. 12 versus the radial separation r. It decreases exponentially as a function of the separation between the antisites, showing the existence of a repulsive interaction between the antisites.

At this stage we discuss the origin of the repulsion. We fitted different Coulomb interaction models to the data. The best fit is shown as blue line in Fig. 12. It is a screened Coulomb potential for three-dimensional bulk screening, modified by a surface dielectric constant. This result indicates that the antisites are charged near the growth surface, during their incorporation. At the GaAs (001) surface, the Fermi level is known to be pinned mid gap [65]. Figure 8 indicates that the (+/0) charge transition level is located above the Fermi energy, when the Fermi level is positioned at midgap. Hence, the antisite defects are positively charged during incorporation in the LTG layer at the growth surface. With progressing growth and hence increasing distance from the growth surface, the conduction band approaches the Fermi level, due to the use of n-type GaAs substrate, and thereby the antisites become uncharged. This neutral charge state is consistent with the result obtained by EPR. The repulsive interaction that is seen in the STM images is thus reminiscent of the positive charge state of the antisites at the time of the growth. But during their observation in STM images the antisites are negatively charged due to the proximity of the STM tip.

VI. CONCLUSION

In summary, we have used complimentary techniques to identify the nature of point defects incorporated in low-temperature grown dilute phosphide GaAs1-xPx alloys. We show that the excess of group V elements gives rise to a preferential incorporation of AsGa instead of PGa, while P is incorporated substitutional on anion sites. Based on the STM analysis of the LTG layers, we identify three different charge states for the AsGa antisites: they are positively charged at the (001) growth surface, uncharged in the bulk LTG layer and negatively charged at the (110) surface in the presence of positively biased STM tip. Furthermore, we observe a tip-induced bond rupture in the direct vicinity of the antisites at two specific lattice positions creating As vacancy-AsGa antisite complexes. The weak bond strength at these positions is in contrast to antisites in pure LTG GaAs and could be related to the presence of P impurities.

Despite a change of the composition from pure GaAs to a dilute phosphide GaAs1-xPx alloy, the critical physical properties of the AsGa antisites in the GaAs1-xPx matrix are kept identical: they give rise to subpicosecond carrier lifetimes and they condensate into precipitates upon annealing. Hence, LTG GaAs1-xPx alloys are even preferable, since they reduce the strain in the epilayer, while keeping the same amount of excess arsenic in comparison with pure LTG GaAs. Thus, LTG GaAs1-xPx offers the advantage to reach higher critical thicknesses. This result also shows that the growth of a Ga51In49P buffer layers in device fabrication is not detrimental to the electronic properties of a subsequently LTG GaAs contaminated with residual phosphorus atoms.

  order to identify these defects, we first analyze the electrical nature of the LTG layer at the local scale. The comparison of tunneling current spectra measured on the n-type GaAs (110) surface and on well-ordered areas of the LTG (110) surface shows a different electrical characteristic between both layers [Fig. 2(b)]. While a zero-current region extending over 1.5 V is measured on the n-type surface, with the Fermi level positioned close to the conduction band edge (EC), we observe a larger apparent band gap on the LTG layer, with EC shifted towards higher energies. The computation of the tunneling current for the n-type GaAs (110) surface shows good agreement with the experimental curve for a tip sample separation of 0.55 nm, a tip radius of 40 nm and a tip work function of 5.6 eV. In addition, we

  predominant row orientation rotated by 90° from [001] to oriented rows. Such a depression reveals that the point defects are negatively charged. A negative charge center induces a local upward band bending, which prevents the tip-induced formation of the accumulation layer in the conduction band. In the constant-current mode, this effect results in a decrease of the tunneling current, due to the localized absence of electrons in the conduction band, compensated by a reduction of the tip-to-sample distance to keep the current constant. Only the filled valence band states contribute to the tunneling current around the defects, accounting for a change of the row orientation. The observed vertical rows correspond to the A5 dangling bond state localized at the anions (imaging of the anion sublattice). Based on the identification of the atomic rows seen in the surrounding of the point defects, we now analyze the symmetry of the defect-related features. Figure 3 shows high resolution STM images of three major types of defects that are found in Fig. 2(a). In Fig. 3(a), the type A defect consists of three bright protrusions, a central one and two others pointing in and directions. It is cut by a mirror plane, shown as two dotted segments, that runs between two Ga rows. The central protrusion is thus located on a surface As sublattice row along the [001] direction. Moreover, the center of the triangle defined by the three protrusions is located between two surface As sublattice rows running along the direction. Such a position is consistent with a Ga site in the second atomic plane, taking the first plane as the surface (see Fig. 3(b)). Therefore, we believe that this point defect is an anion antisite belonging to the second atomic plane. This is corroborated by the similarity of the contrast with the one theoretically predicted[34,35,36], despite the smaller size of its central core.While the comparison of the atomic contrast for type B and C defects visible in Fig.3(b) and (c) reveals some similarities with a type A defect, both defects distinguish themselves by the contrast of their satellites. The satellites consist of either one protrusion and one atomic-sized depression (type B defect) or two atomic-sized depressions (type C defect). For the type B defect, the protrusion and the depression can be found inverted with respect to the mirror plane. From the subsequent acquisition of sequential STM frames such as the one shown in Fig.3f, it is clear that the defects randomly switch from one type of defect to another type, sometimes with some intermediate configurations. For example, the central type B defect in Fig.3f1switches to a type B' defect (f2→f3), that turns into a type C defect (f3→f4). Then the type C defects switches to a type A defect (f4→f5). Based on these observations, we conclude that the defects seen in Fig.3(b) and (c) correspond to an anion antisite in the second atomic plane and two anion third-neighbor surface As atoms in the and directions that show strong instabilities. Counting the number of all these anion antisites, we find a corresponding bulk concentration of 7.1±0.2 x 10 19 cm -3 in the LTG layer grown at 240°C, consistent with the high concentration of antisites usually encountered in LTG GaAs grown with the same temperature[3,37,38].

Fig. 2a .

 2a Fig. 2a. Based on the isosurfaces of the electron density for an antisite positioned in the second layer, it appears that a negative charge implies the partial filling of the closest Ga dangling bonds in the and directions. This change of the electron distribution might weaken the As-Ga bonds at the surface. When the tip scans above the As surface atoms, the subsequent electronic excitation of the weakened bonds might favor the desorption of the atoms, through a two hole localization mechanism for example [62,63], leaving a vacancy behind. If an As adatom stick to the tip apex instead of binding with

FIG. 1 .FIG. 3 .FIG. 6 .FIG. 7 .FIG. 8 .FIG. 9 .FIG. 11 .FIG. 12 .

 1367891112 FIG. 1. SIMS concentration depth profiles of As, Ga, P, and In for a sample with a LTG layer grown at 240°C. Inset: Cross-sectional HAADF-STEM image showing the basic structure of the sample.

  FIGURE 4

Figure 12

 12 Figure 12

ACKNOWLEDGMENTS

This study was financially supported by the DGA (Direction Générale de l'Armement), the RENATECH network, the French National Research Agency (the Equipex Programs Excelsior ANR-

11-EQPX-0015), the Agence Universitaire de la Francophonie (Grant N° BAO-2012-05-52220FT2E0).

The authors thank C. Delerue for useful discussions. * Corresponding author: bruno.grandidier@isen.iemn.univ-lille1.fr