Computational and experimental imaging of Mn defects on GaAs (110)
  cross-sectional surface
Computational and experimental imaging of Mn defects
on GaAs (110) cross-sectional surface
A. Stroppa,1, 2, ∗ X. Duan,1, 2, † M. Peressi,1, 2, ‡ D. Furlanetto,3, 4 and S. Modesti3, 4
1Dipartimento di Fisica Teorica, Università di Trieste,
Strada Costiera 11, 34014 Trieste, Italy
2CNR-INFM DEMOCRITOS National Simulation Center,
via Beirut 2-4, 34014 Trieste, Italy
3CNR-INFM TASC National Laboratory,
Area Science Park, 34012 Trieste, Italy
4Dipartimento di Fisica and Center of Excellence for Nanostructured Materials, CENMAT,
Università di Trieste, via A. Valerio 2, 34127 Trieste, Italy
(Dated: August 10, 2021)
Abstract
We present a combined experimental and computational study of the (110) cross-sectional surface
of Mn δ-doped GaAs samples. We focus our study on three different selected Mn defect configu-
rations not previously studied in details, namely surface interstitial Mn, isolated and in pairs, and
substitutional Mn atoms on cationic sites (MnGa) in the first subsurface layer. The sensitivity of
the STM images to the specific local environment allows to distinguish between Mn interstitials
with nearest neighbor As atoms (IntAs) rather than Ga atoms (IntGa), and to identify the finger-
print of peculiar satellite features around subsurface substitutional Mn. The simulated STM maps
for IntAs, both isolated and in pairs, and MnGa in the first subsurface layer are consistent with
some experimental images hitherto not fully characterized.
PACS numbers: 73.20.-r,73.43.Cd,68.37.Ef
http://arxiv.org/abs/0704.1357v1
I. INTRODUCTION
Mn-doped GaAs1,2,3,4 has attracted considerable attention among the diluted magnetic
semiconductors for its possible application in the emerging field of spintronic.5,6,7 Although
other materials such as ferromagnetic metals and alloys, Heusler alloys, or magnetic oxides
seem to be promising candidates for spintronic devices, the diluted magnetic semiconduc-
tors and Mn-doped GaAs in particular are of tremendous interest in that they combine
magnetic and semiconducting properties and allow an easy integration with the well es-
tablished semiconductor technology. Besides possible spintronic applications, characterizing
and understanding the properties of Mn defects in GaAs is a basic research problem which
is still debated.
The growth conditions and techniques affect the solubility of Mn in GaAs, which is in
general rather limited, and its particular defect configurations, thus determining the mag-
netic properties of the samples.8,9,10,11,12 The highest Curie temperature Tc reachable for
Mn-doped GaAs up to few years ago was 110 K,13 rather low for practical technological
purposes. Intense efforts have been pursued in the last years in order to understand the
physics of this material and to improve its quality and efficiency. Out-equilibrium growth
techniques1,5 have enabled to increase the solubility of Mn and the Curie temperature;
post-growth annealing of epitaxial samples at temperatures only slightly above the growth
temperature has been particularly successfull.9,10,14 Nowadays, δ-doping is used as an al-
ternative to the growth of bulk MnxGa1−xAs,
15,16 allowing to obtain locally high dopant
concentrations and, remarkably, an important enhancement of Tc, up to about 250 K.
17,18
For further improvements it is essential to investigate the different configurations of Mn
impurities and their effect on the magnetic properties of the system. The most common
and widely studied Mn configuration is substitutional in the cation sites (MnGa), with Mn
acting as a hole-producing acceptor.17 To a less extent, Mn can also occupy interstitial sites,
in particular tetrahedral ones. In such a case, it is expected to strongly modify the magnetic
properties, acting as an electron-producing donor and hence destroying the free holes and
hindering ferromagnetism.19
Interstitials have not been fully characterized up to now, although their existence has been
suggested in different situations.8,9,10,11,14,20,21,22,23,24,25,26 For instance, the enhancement of
the Curie temperature after post-growth annealing has been attributed to the reduction of
interstitial defects with their out diffusion towards the surface.11 It has been suggested that
interstitial sites are highly mobile and could be immobilized when adjacent to substitutional
MnGa, thus forming compensated pairs with antiferromagnetic coupling.
27 A first identifi-
cation of interstitial Mn dates back to almost fifteen years ago by electron paramagnetic
resonance (EPR).20 Very recently EPR spectra from variously doped and grown samples of
Mn-doped epitaxial GaAs have allowed to identify the presence of ionized Mn interstitials
at concentrations as low as 0.5%, although not providing details about the specific local en-
vironment of the interstitial site.28 Recent X-ray absorption near edge structure (XANES)
and extended x-ray absorption fine structure (EXAFS) spectra in Mn δ-doped GaAs samples
suggest that Mn occupy not only substitutional Ga sites but also interstitial sites, mainly
in case of Be co-doping.29
Cross-sectional Scanning Tunneling Microscopy (XSTM) allows a direct imaging of the
electronic states and can be used to characterize the impurities near the cleavage surface.30
In recent years several XSTM studies of Mn-doped GaAs samples have been performed but
without a complete consensus on the defects characterization.22,31,32,33,34,35,36,37,38 We stress
that most of XSTM studies mainly concern MnxGa1−xAs alloys and have identified mainly
substitutional Mn defects. δ-doped samples have been investigated by Yakunin et al.,35 who
pointed out the advantage that in such samples it is easy to discriminate Mn related defects
from other defects.
From the theoretical point of view, numerical works have been also focused mainly on the
simulation of XSTM images of substitutional impurities on uppermost surface layers.31,33,34,35
A complete and detailed investigation of interstitial impurities as they can appear on the
exposed cleaved surface is still lacking thus preventing the possibility of a comprehensive
interpretation of all the available experimental XSTM images.
Mn δ-doped (001) GaAs samples recently grown at TASC Laboratory in Trieste and
analyzed with XSTM on the (110) cleavage surface have shown several Mn related features
(see Fig. 1). Some of them have already be studied by other groups, like the asymmetric
cross-like (or butterfly-like) structures marked by A in Fig. 1(a), attributed to Mn acceptors
a few atomic layers below the surface.34 Some other features, such those marked by B, or
those of Fig 1(b), have not been yet assigned to specific Mn configurations. In order to
identify the kind of Mn defects that cause them we have performed new density functional
simulation of cross-sectional XSTM images focusing on three selected defect configurations
not yet fully studied, but whose presence cannot be excluded in real samples. In particular,
we focus our attention on interstitial surface configurations, both individual as well as in
pairs. We have also considered MnGa on the first layer below the surface and compared all
the simulations with the experimental maps.
II. EXPERIMENTAL DETAILS
Mn δ-doped samples were grown by molecular beam epitaxy on GaAs(001) in a facility
which includes a growth chamber for III-V materials and a metallization chamber. After the
growth of a Be doped buffer at 590oC and of an undoped GaAs layer 50 nm thick at 450oC
with an As/Ga beam pressure ratio of 15, the samples were transferred in the metallization
chamber where a submonolayer-thick Mn layer was deposited at room temperature at the
rate of 0.003 monolayer/s. An undoped GaAs cap layer was subsequently grown at 450oC.
This procedure was repeated in order to have three δ-doped Mn layer of 0.01, 0.05 and 0.2
monolayers in the same sample. During the transfers and the Mn deposition the vacuum
was always better than 2×10−8 Pa. The 0.1 mm thick wafers containing the Mn layers were
cleaved in situ in a ultra high vacuum STM system immediately prior to image acquisition
to yield atomically flat, electronically unpinned {110} surfaces containing the [001] growth
direction and the cross section of the δ-doped layer. The XSTM image presented in Fig. 1
and the others shown in this paper have been acquired from a δ-doped Mn layer of 0.2
monolayers with W tips.
The densities of the features observed by XSTM near each Mn layer were approximately
proportional to the Mn coverage of the δ-doped layer in the range 0.01-0.2 monolayer. No
trace of contaminants was observed by in situ x-ray photoemission spectroscopy after the
transfer in the metallization, after the Mn deposition, and after the transfer in the growth
chamber. For these two reasons we attribute the features observed by XSTM to the Mn
atoms, and not to defects or contaminants caused by the growth interruption and transfers
between the chambers. The density of the defects caused by these steps should not depend
on the Mn coverage, contrary to what we observe. Moreover, a sample was grown with the
same procedure described above, including the transfers between the chambers, but without
the Mn deposition. The photoluminescence spectra of this sample are undistinguishable from
that of a good undoped GaAs epitaxial layer grown without transfers between the chambers.
This confirms that the transfers do not introduce an appreciable amount of defects.
III. THEORETICAL APPROACH
Our numerical approach is based on spin-resolved Density Functional Theory (DFT) using
the ab-initio pseudopotential plane-wave method PWscf code of the Quantum ESPRESSO
distribution.39 Cross-sectional surfaces are studied using supercells with slab geometries,
according to a scheme previously used,40 with 5 atomic layers and a vacuum region equivalent
to 8 atomic layers. Mn dopants are on one surface, whereas the other is passivated with
hydrogen. For a single Mn impurity we use a 4×4 in-plane periodicity corresponding to
distances between the Mn atom and its periodic images of 15.7 Å along the [11̄0] and 22.2
Å along [001]. No substantial changes in the XSTM images have been found using a 6×4
periodicity, which has been instead routinely used when considering interstitial complexes.
Other details on technicalities can be found in Ref. 41.
In our study, we have mainly focused on the Local Spin Density Approximation (LSDA)
for the exchange-correlation functional. An ultrasoft pseudopotential is used for Mn atom,
considering semicore 3p and 3s states kept in the valence shell while norm-conserving pseu-
dopotentials have been considered for Ga and As atoms. The 3d-Ga electrons are considered
as part of core states.42,43
Tests beyond LSDA (with Generalized Gradient Correction and LSDA+U methods) have
not shown any substantial difference in the features of the XSTM maps. As a further check,
we have also simulated ionized substitutional MnGa (with charge state equal to 1−) on
surface and in the first subsurface layer as well as ionized interstitial Mn (with charge state
equal to 2+) on surface layer. Neither the former nor the latter simulated XSTM maps show
significant differences with respect to the neutral cases. We address the reader to a future
pubblication for details.44
The XSTM images are simulated using the model of Tersoff-Hamann,45,46 where the
tunneling current is proportional to the Local Density of States (LDOS) at the position of
the tip, integrated in the energy range between the Fermi energy Ef and Ef + eVb, where
Vb is the bias applied to the sample with respect to the tip. The position of the Fermi level
is relevant for the XSTM images. In general, Ef strongly depends on the concentration of
dopants: this is contrivedly large in our simulations even in the case of a single Mn dopant
per supercell. Therefore to overcome this problem we fix Ef according to the experimental
indications: in order to account for the p-doping in the real samples, we set Ef close to
the Valence Band Maximum (VBM). The VBM in the DOS of the Mn-doped GaAs can be
exactly identified by aligning the DOS projected onto surface atoms far from the impurity
with the one of the clean surface. In any case, the comparison between experiments and
simulations must be taken with some caution, due to the possible differences in the details
entering in the determination of the XSTM image, such as tip-surface separation, precise
value of the bias voltage and position of Ef , surface band gap.
IV. SURFACE MN INTERSTITIALS
We first focus on interstitial dopant configurations, IntAs and IntGa. Throughout this
work we have considered only tetrahedral interstitial position, since it is known from bulk
calculations that the total energy corresponding to the hexagonal interstitial site is higher
by more than 0.5 eV.11,48,49 The tetrahedral interstitial site in the ideal geometry has four
nearest-neighbor (NN) atoms at a distance equal to the ideal host bond length d1 and six
next-nearest-neighbor (NNN) atoms at the distance d2 =
d1, which are As(Ga) atoms for
IntGa(As), respectively. At the ideal truncated (110) surface, the numbers of NNs and NNNs
reduce to three (2 surface atoms and 1 subsurface atom) and four (2 surface atoms and 2
subsurface atoms) instead of four and six respectively.
In the uppermost panels of Fig. 2 we show a ball and stick side and top view of the
relaxed IntAs and IntGa configurations. In the relaxed structure, due to symmetry breaking
because of the surface and the consequent buckling of the outermost surface layers, the NN
and NNN bond lengths are no longer equal. Furthermore, some relaxed NNs bond lengths
turn out to be longer than NNNs ones. In the following, we do not distinguish among NN
and NNN atoms: they are simply referred as neighbor surface or subsurface atoms, as shown
in the Figure.
The two relaxed configurations slightly differ in energy, by ∼ 130 meV/Mn atom, in
favour of IntGa. This is at variance with the bulk case studied in the literature, where it
has been found that IntAs is favoured: for neutral state, the energy difference is actually so
small (5 meV/Mn atom)48 that it is not meaningful, but it goes up to 350 meV in case of
interstitial Mn with 2+ charge state.11
After optimization of the atomic positions, sizeable displacements from the ideal zinc
blende positions occur for the Mn impurities and their surface and subsurface neighbors;
small relaxations effects are still present in the third layer, in both configurations. In IntAs,
with respect to the ideal (110) surface plane, Mn relaxes outward by ∼ 0.06 Å and Assurf
(Assubsurf) move upwards (downwards). On the other hand, the Ga atoms (both on surface
and subsurface) are shifted towards the bulk. In IntGa, Mn relaxes inward by ∼ 0.32 Å; the
Gasurf and Gasubsurf atoms are displaced downwards and the Assurf (Assubsurf) atom moves
upwards (downwards). The interatomic distances between Mn and the nearest atoms are in
general longer by more than 2-3 % than ideal values (details in Ref. 41).
The simulated XSTM images of IntAs (left) and IntGa (right) configurations at negative
and positive bias voltages (from − 2.0 V to +2.0 V) are shown in the lower panels of Fig. 2.
In IntAs, Mn appears as an additional bright spot at negative bias voltage (Vb=−1 V),
slightly elongated in the [001] direction and located near the center of the surface unit cell
identified by surface As atoms. The Assurf atoms close to Mn appear less bright than the
others. These features are similar changing Vb from −1 to −2 V.
In the empty states image at Vb=1 V Mn appears again as an elongated bright spot. The
underlying cation lattice is only barely visible at this bias voltage. The very bright XSTM
feature originates from the Mn d minority states and a strong peak of Gasurf majority
states.50 At Vb=2 V, this feature is still well visible, as well as another region brighter
than the underlying cationic sublattice in correspondence of Assurf atoms neighbor to Mn,
suggesting a contribution coming from the hybridization between Mn-d and Assurf -p states.
In IntGa configuration, at negative voltage, Mn appears as an almost circular bright spot
located in between two surface As atoms adjacent along the [001] direction. At positive bias
voltages, the two Gasurf atoms neighbor to Mn appear very bright with features extending
towards Mn in a “v”-shaped form and the atoms in the neighborhood also look brighter
than normal. These features remain visible by increasing the positive bias voltage up to 2
eV. Remarkably the empty states images of Mn are quite different for the two interstitial
configurations, making them clearly distinguishable by XSTM analysis. Some features in
the experimental XSTM images appear as bright spots both at positive and negative bias
voltages. These spots lie along the [001] Ga rows and between the [1-10] Ga columns at
positive bias voltage (see Fig. 1(b)). Their location with respect to the surface Ga lattice
and the comparison with the simulated images allow to identify them as IntAs Mn atoms.
The numerical simulation gives easily informations on the magnetic properties of the
system. The total and absolute magnetization, calculated from the spatial integration of
the difference and the absolute difference respectively between the majority and minority
electronic charge distribution, are different in the two configurations: 4.23 and 4.84 µB
for IntAs and 3.41 and 4.71 µB for IntGa respectively. These differences indicate in both
cases the presence of region of negative spin-density and a clear dependence of the induced
magnetization on the local Mn environment. The individual atomic magnetic moments
can be calculated as the difference between the majority and minority atomic-projected
charges. In IntAs, Mn magnetic moment is 3.96 µB, almost integer, corresponding to the
presence of a gap in the Mn-projected minority density of states. Mn magnetization is
slightly lower in IntGa (3.67 µB). In both cases they are significantly larger compared to
the bulk case, indicating a surface induced enhancement. The analysis of spin-polarization
induced by interstitial Mn on its nearest neighbors shows in both cases an antiferromagnetic
Mn–Ga coupling and a smaller ferromagnetic Mn–As coupling: more precisely, the magnetic
moments induced on surface Ga atoms neighbors to Mn are equal to −0.14 and −0.17 µB
in IntAs and IntGa respectively, whereas those induced on surface or subsurface As atoms
neighbors to Mn are positive and at most equal to 0.05 µB. We address the reader to Ref.
for further details.
In the experimental images of Mn δ-doped GaAs samples we often observe two spots
close one each other at a distance of about 8 Å, as reported in Fig. 3 (larger panel). The
simulated image of two IntAs atoms separated by a clean surface unit cell along (110),
partially superimposed, reproduces the main features of this experimental image, and it is
basically a superposition of images of individual IntAs (elongated bright spot each one, with
major axis along the [001] direction, and a surrounding darker region).
V. SUBSTITUTIONAL MN DEFECTS IN THE FIRST SUBSURFACE LAYER
Another typical feature present in the experimental XSTM maps is a bright spot visible
at positive bias voltages with two satellite features forming a triangular structure, as shown
in Fig. 1(a) (feature B) and in Fig. 4 in the lower panels. This feature seems similar to
that caused by the arsenic antisite defect (As on Ga) in GaAs.51,52 However in the arsenic
antisite defects the satellites are visible only at negative sample bias, while the defect that we
observe in the Mn layers shows satellite only in the positive bias images. On the other hand
there is a clear resemblance of the defect B (Fig. 1(a) and Fig. 4) with the simulated image
of a substitutional MnGa atom in the first subsurface layer shown in the panels partially
superimposed to the experimental images. It can be seen at Vb < 0 a deformation of the
surface As rows in correspondence of the Mn impurity below, and, even more remarkably,
the peculiar satellite bright features on two neighboring surface As stoms at Vb > 0 giving
rise to a triangular-shaped image. Therefore we attribute the defect B to substitutional Mn
Ga atoms in the first subsurface layer.
Finally, we discuss our findings in comparison with some relevant results present in the
literature. The comparison of our simulations with those of Sullivan et al.33 is possible only
for the isolated MnGa in the first subsurface layer at negative bias voltage: in such a case
the simulated images show similar features. The corresponding image at positive bias is not
reported and other configurations are not comparable.
The XSTM imaging of substitutional Mn is reported with more details by Mikkelsen et
al.,31,32 where both the simulated maps for surface and subsurface MnGa and the experimental
ones attributed to this impurity configuration are shown at negative and positive bias, thus
allowing for a more complete comparison. The images for MnGa in the first subsurface layer
have a good resemblance with ours, a part from the satellite features that we have identified
at positive bias on neighbor As atoms which are not present in their images, neither in
the simulated nor in the experimental one. More precisely, we notice that their simulated
surface area is too small to make such satellite features visible. The simulated images for
surface MnGa are also similar to ours and, like ours, not corresponding to any experimental
feature.44 This leads to the conclusion that the presence of substitutional Mn in the first
layer of the exposed surface is very unlikely.
Mikkelsen et al. reported also the simulation of surface interstitial Mn in their Fig.
3(d),32 that according to our understanding on the basis of the symmetry planes should
correspond to IntGa, although not explicitely indicated. Their images are similar to ours for
the same configuration. They rule out the presence of interstitials since these images are
not compatible with experiments, at variance with our findings concerning IntAs. It should
be noted however that we observe the IntAs features in the experimental samples only in the
first few hours after the sample cleavage. They disappear for longer times, probably because
of surface contamination or diffusion.
Kitchen et al.36,37 report experimental images for Mn adatoms at the GaAs (110) surface
with highly anisotropic extended star-like feature, attributed to a single surface Mn acceptor.
Interestingly, these images are compatible with our simulated surface MnGa, not show here.
A resemblance with our empty state image for IntAs (see Fig. 2 at Vb=+2 V) is instead only
apparent because the mirror symmetry plane is different.
An anisotropic, crosslike feature in XSTM image is reported also Yakunin et al.34 and,
from comparison with an envelope-function, effective mass model and a tight-binding model,
it is attributed to a hole bound to an individual Mn acceptor lying well below the surface.
We observe similar feature of different sizes (see Fig. 1), the smallest of them are those
reported in Fig. 4, that we identify as MnGa in the first subsurface layer.
A part from different details, our simulated images for surface and subsurface MnGa
are compatible with such crosslike features, although experimental and simulated images
reported therein concern substitutional impurities located more deeply subsurface than those
we have considered. Crosslike features are observed even at very short Mn-Mn spatial
separations.35
VI. CONCLUSIONS
We have reported a combined experimental and first-principles numerical study of XSTM
images of the (110) cross-sectional surfaces of Mn δ-doped GaAs samples. We suggest
an identification of three typical configurations observed in the experimental sample on
the basis of a comparison of numerical prediction and observed images both at negative
and positive applied bias. (i) Some structures observed can be identified as surface Mn
interstitial with As nearest neighbors, on the basis of their position with respect to the
surface lattice and the comparison with the simulated images. At variance, there is no
evidence in the experimental samples of Mn interstitial with Ga nearest neighbors, whose
XSTM imaging according to our numerical simulations would correspond to very different
features. (ii) Besides isolated configurations, also pairs of Mn interstitials with As nearest
neighbors are clearly observed and identified. (iii) Subsurface substitutional MnGa atoms in
the first subsurface layer can also be unambigously identified in the experimental images by
a main bright spot corresponding to the dopant and from peculiar satellite features on two
neighboring As atoms which are clearly observed in the experimental images and predicted
by simulations.
VII. ACKNOWLEDGMENTS
Computational resources have been partly obtained within the “Iniziativa Trasversale di
Calcolo Parallelo” of the Italian CNR-Istituto Nazionale per la Fisica della Materia (CNR-
INFM) and partly within the agreement between the University of Trieste and the Consorzio
Interuniversitario CINECA (Italy). We thank A. Franciosi, S. Rubini and coworkers for the
preparation of the sample and fruitful comments and discussions; A. Debernardi for his help
in the pseudopotential generation and for useful discussions. Ball and stick models and
simulated images are obtained with the package XCrySDen.53
∗ Electronic address: astroppa@ts.infn.it; Presently at: Institute of Material Physics, University
of Vienna, Sensengasse 8/12, A-1090 Wien, Austria and Center for Computational Materials
Science (CMS), Wien, Austria
† Presently at School of Physics, The University of Sydney, NSW 2006 Australia
‡ Electronic address: peressi@ts.infn.it
1 K. Takamura et al., J. Appl. Phys. 89, 7024 (2001).
2 H. Ohno, A. Shen, F. Matsukura, A. Oiwa, A. Endo, S. Katsumoto, and Y. Iye, Appl. Phys.
Lett. 82, 3020 (2003).
3 H. Ohno, F. Matsukura, and Y. Ohno, Mater. Sci. Eng. B 84, 70 (2001).
4 T. Jungwirth, Jairo Sinova, J. Mašek, J. Kučera and A.H. MacDonald, Rev. Mod. Phys. 78,
809 (2006).
5 H. Ohno, Science 281, 51 (1998), and references therein.
6 Y. Ohno, D.K. Young, B. Beschoten, F. Matsukura, H. Ohno, and D.D. Awschalom, Nature
402, 790 (1999).
7 H. Ohno, D. Chiba, F. Matsukura, T. Omiya, E. Abe, T. Dietl, Y. Ohno, and K. Ohtani, Nature
408, 944 (2000).
8 K.M. Yu, W. Walukiewicz, T. Wojtowicz, I. Kuryliszyn, X. Liu, Y. Sasaki, J.K. Furdyna, Phys.
Rev. B 65, 201303(R) (2002).
9 K. W. Edmonds, K. Y. Wang, R. P. Campion, A. C. Neumann, N. R. S., Rarley, B. L. Gallagher,
and C. T. Foxon, Appl. Phys. Lett. 81, 4991 (2002).
10 K. C. Ku, S. J. Potashnik, R. F. Wang, S. H. Chun, P. Schiffer, N. Samarth, M. J. Seong, A.
Mascarenhas, E. Johnston-Halperin, R. C. Mayers, A. C. Gossard, and D. D. Awschalom, Appl.
Phys. Lett. 82, 2302 (2003).
11 K.W. Edmonds, P. Boguslawski, K.Y. Wang, R.P. Campion, S.N. Novikov, N.R.S. Farley, B.L.
Gallagher, C.T. Foxon, M. Sawicki, T. Dietl, M.B. Nardelli, and J. Bernholc, Phys. Rev. Lett.
92, 037201 (2004).
12 L. Bergqvist, P. A. Korzhavyi, B. Sanyal, S. Mirbt, I. A. Abrikosov, L. Nordström, E.A.
Smirnova, P. Mohn, P. Svedlindh, and O. Eriksson, Phys. Rev. B 67, 205201 (2003).
13 F. Matsukura, H. Ohno, A. Shen, Y. Sugawara, Phys. Rev. B 57, R2037 (1998).
mailto:astroppa@ts.infn.it
mailto:peressi@ts.infn.it
14 T. Jungwirth, K. Y. Wang, J. Mašek, K. W. Edmonds, Jürgen König, Jairo Sinova, M. Polini,
N. A. Goncharuk, A. H. MacDonald, M. Sawicki, A. W. Rushforth, R. P. Campion, L. X. Zhao,
C. T. Foxon, and B. L. Gallagher, Phys. Rev. B 72, 165204 (2005).
15 A.M. Nazmul, S. Sugahara, and M. Tanaka, Phys. Rev. B 67, 241308(R) (2003).
16 E. F. Schubert, J. M. Kuo, R. F. Kopf, H. S. Luftman, L. C. Hopkins, and N. J. Sauer, J. Appl.
Phys. 67, 1969 (1990).
17 T. Dietl, H. Ohno, F. Matsukura, J. Cibert, and D. Ferrand, Science 287, 1019 (2000).
18 A.M. Nazmul, T. Amemiya, Y. Shuto, S. Sugahara, and M. Tanaka, Phys. Rev. Lett. 95, 017201
(2005).
19 P. Mahadevan and A. Zunger, Phys. Rev. B 68, 075202 (2003).
20 S.J.C.H.M. van Gisbergen, M. Godlewski, T. Gregorkiewicz, and C.A.J. Ammerlaan, Appl.
Surf. Sci. 50, 273 (1991).
21 F. Glas, G. Patriarche, L. Largeau, and A. Lemaitre, Phys. Rev. Lett. 93, 086107 (2004).
22 G. Mahieu, P. Condette, B. Grandidier, J. P. Nys, G. Allan, D. Stivenard, Ph. Ebert, H.
Shimizu, and M. Tanaka, Appl. Phys. Lett. 82, 712 (2003).
23 K. W. Edmonds, N. R. S. Farley, T. K. Johal, G. van der Laan, R. P. Campion, B. L. Gallagher,
and C. T. Foxon, Phys. Rev. B 71, 064418 (2005).
24 R. Wu, Phys. Rev. Lett. 94, 207201 (2005).
25 S.C. Erwin and A.G. Petukhov, Phys. Rev. Lett. 89, 227201 (2002).
26 V. Holý, Z. Matěj, O. Pacherová, V. Novák, M. Cukr, K. Olejnik, and T. Jungwirth, Phys. Rev.
B 74, 245205 (2006).
27 J. Blinowski, P. Kacman, Phys. Rev. B 67, 121204(R) (2003).
28 T. Weiers, Phys. Rev. B 73, 033201 (2006).
29 F. d’Acapito, G. Smolentsev, F. Boscherini, M. Piccin, G. Bais, S. Rubini, F. Martelli, and A.
Franciosi, Phys. Rev. B 73, 035314 (2006).
30 R. M. Feenstra, Semicond. Sci. Technol. 9, 2157 (1994).
31 A. Mikkelsen, B. Sanyal, J. Sadowski, L. Ouattara, J. Kanski, S. Mirbt, O. Eriksson, and E.
Lundgren, Phys. Rev. B 70, 85411 (2004).
32 A. Mikkelsen, E. Lundgren, Progress in Surf. Sci. 80, 1 (2005).
33 J. M. Sullivan, G. I. Boishin, L. J. Whitman, A. T. Hanbicki, B. T. Jonker, and S. C. Erwin,
Phys. Rev. B 68, 235324 (2003).
34 A.M. Yakunin, A.Y. Silov, P.M. Koenraad, J.H. Wolter, W. Van Roy, J. De Boeck, J.M. Tang,
M.E. Flatté, Phys. Rev. Lett. 92, 216806 (2004).
35 A.M. Yakunin, A.Yu. Silov, P.M. Koenraad, J.-M. Tang, M.E. Flatté, W. Van Roy, J. De Boeck,
J.H. Wolter, Phys. Rev. Lett. 95, 256402 (2005).
36 D. Kitchen, A. Richardella, A. Yazdani, J. Supercond. 18, 23 (2005).
37 D. Kitchen, A. Richardella, J. Tang, M.E. Flattè, and A. Yazdani, Nature 442, 436 (2006).
38 J.N. Gleason, M.E. Hjelmstad, V.D. Dasika, R.S. Goldman, S. Fathpour, S. Charkrabarti, and
P.K. Bhattacharya, Appl. Phys. Lett. 86, 011911 (2005).
39 http://www.pwscf.org and http://www.quantum-espresso.org
40 X. Duan, M. Peressi, and S. Baroni, Phys. Rev. B 72, 085341 (2005); X. Duan, S. Baroni, S.
Modesti, and M. Peressi, Appl. Phys. Lett. 88, 022114 (2006).
41 A. Stroppa and M. Peressi, Mat. Sci. Eng. B 126 217 (2006).
42 Norm-conserving pseudopotentials from the publicly available Quantum ESPRESSO table are
used: As.pz-bhs.UPF, Ga.pz-bhs.UPF, H.pz-vbc.UPF, and the pseudopotential for Mn used in
Ref. 43
43 A. Debernardi, M. Peressi, and A. Baldereschi, Mat. Sc. Eng. C 23, 743 (2003) and other more
recent works.
44 A. Stroppa and M. Peressi, to be published.
45 J. Tersoff and D.R. Hamann, Phys. Rev. Lett. 50, 1998 (1983).
46 J. Tersoff and D.R. Hamann, Phys. Rev. B 31, 805 (1985).
47 However, we have checked that a variation of ∼ ± 0.3 eV in the bias voltage considered in the
simulated images does not affect their basic features.
48 J. Mašek and F. Máca, Phys. Rev. B 69, 165212 (2004).
49 J.X. Cao, X.G. Gong, and R.Q. Wu, Phys. Rev. B 72, 153410 (2005).
50 A. Stroppa, Nuovo Cim. 29, 315 (2006).
51 R.M. Feenstra, J.M. Woodall and G.D. Pettit, Phys. Rev. Lett. 71, 1176 (1993).
52 G. Mathieu et al., Appl. Phys. Letters 82, 712 (2003).
53 A. Kokalj, Comp. Mater. Sci., 2003, Vol. 28, p. 155. Code available from
http://www.xcrysden.org/.
http://www.pwscf.org
http://www.xcrysden.org/
FIG. 1: (a) Experimental (110) XSTM image of a 0.2 monolayer Mn δ-doped layer in GaAs at a
sample bias voltage of 1.7 eV. This image has not been corrected for the drift of the sample. (b)
XSTM image of a Mn related structure at the bias voltage of −1.4 eV (left) and +1.9 eV (right).
The white lines show the [001] Ga atomic rows.
FIG. 2: Isolated Mn interstitial dopants on GaAs(110) surface, with As nearest neighbors (IntAs,
left) and Ga nearest neighbors (IntGa, right). Upper panels: ball-and-stick model of the relaxed
surface, top and side view. Only the three topmost layers are shown in the side view. Black spheres
are Mn, white spheres are As, grey spheres are Ga. Lower panels: simulated XSTM images at
occupied states and empty states respectively, for different bias voltages.
FIG. 3: Smaller superimposed panel: simulated XSTM image of a pair of IntAs on GaAs(110)
surface with a relative distance of ∼ 8 Å along the [11̄0] direction at a bias voltage Vb=−2 V. The
larger panel shows an experimental image compatible with the simulation.
FIG. 4: Upper smaller superimposed panels simulated XSTM image of a subsurface MnGa on
GaAs(110) at negative (left) and positive (right) bias voltages. The lower panels show correspond-
ing experimental images of the structure B (see Fig. 1(a)) taken at sample bias voltages of −1.4
V (left) and +1.8 V (right) that are compatible with the simulations, performed with voltages of
−1 V and +1 V.
Fig. 1
Fig. 2
Fig. 3
Fig. 4
	Introduction
	Experimental details
	Theoretical approach
	Surface Mn interstitials
	 Substitutional Mn defects in the first subsurface layer
	Conclusions
	Acknowledgments
	References
