# Toward a Detailed Understanding of Si(111)-7×7 Surface and Adsorbed Ge Nanostructures: Fabrications, Structures,   and Calculations

Authors: Ye-Liang Wang; Hai-Ming Guo; Zhi-Hui Qin; Hai-Feng Ma; Hong-Jun Gao
Journal: Journal of Nanomaterials
 (2008)

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## Abstract

Firstly, both the rest atoms and the adatoms of Si(111)-7×7 surface are observed simultaneously by scanning tunneling microscopy (STM) when the sample bias voltages are kept less than − 0.7 V. The visibility of the rest atoms is rationalized by first-principle calculations and a very sharper tip can resolve them. Secondly, the behaviors of various Ge nanostructures fabricated on Si(111)-7×7, ranging from the initial adsorption
sites of individual Ge atoms to the aggregation patterns of Ge
nanoclusters, and then to 2D extended Ge islands, are
comprehensively investigated by STM. The individual Ge atoms
tend to substitute for Si adatoms at
Si(111)-7×7 with the preference of corner adatoms in the faulted half unit  when keeping substrate at 150∘C. With increasing Ge coverage, individual Ge atoms and Ge nanoclusters coexist on the substrate. Subsequently, the density of Ge nanoclusters increase and cluster-distribution becomes gradually regular with the formation of final 2D extended hexagonal configuration. When keeping the substrate at 300∘C, Ge islands consisting of more complicated reconstructions with intermixing Ge/Si components are present on the substrate. The detail structural characterizations and the bonding nature of the observed Ge nanostructures are enunciated by the first-principle calculations.


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## Body

## 1. “Ultimate” STM Images of the Si(111)-7×7 Surface

Si(111)-7× 7 surface as
one of the most complicated and fascinating object of study is being
extensively used in various research fields ranging from surface science and
material science to nanotechnology. As a classic example, this reconstructed
surface provides a platform for the testing of the unprecedented resolution of
STM as a novel powerful apparatus in the early 1980s [1]. The first real space
atomic image of this surface was obtained by Binnig et al. in their
landmark STM experiment [2], in which twelve bright spots corresponding to the
topmost adatoms are revealed.Since then, with this powerful tool and the later
family of scanning probe microscope, the structure of Si(111)-7× 7 surface has
been extensively investigated [3–8]. The demonstrations on the atomic
topography of clean Si(111)-7 × 7 surface
commonly show the topmost adatoms. On the mapping of rest atoms of Si(111)-7 × 7 surface, some
saddle points at the position expected for the rest atoms were reported by
Avouris and Wolkow [3] and Nishikawa et al. [4] using STM. Recently,
some special techniques were used to obtain the images of Si(111)-7 × 7 surface with
atomic scale resolution, such as Lantz et al. [5] using scanning force
microscopy and Giessibl et al. [6, 7] using atomic force microscopy.
Sutter et al. [8] have mapped selectively the rest atoms at a price of
suppressing the adatom spots with a monocrystalline semiconductor tip since its
energy gap can suppress the tunneling from the adatoms at certain sample bias.
STM is very sensitive to states closest to the sample Fermi Energy (EF). The state of dangling bonds
of the adatoms is about 0.4 eV below EF and that of rest atoms is
about 0.8 eV below  EF [9], so it is difficult to map
the rest atoms whose dangling bonds state is far from EF. Up till now, the adatoms and
the rest atoms of Si(111)-7 × 7 surface are
still not clearly distinguished simultaneously by using conventional tungsten
tip. This inability has led to the perception that the measured tunneling
current for semiconductor materials comes mostly
from states near
the Fermi level instead of the states further away, due to the exponential
dependence of the tunneling probability on the energy level position [10]. We revisit
this surface by using STM. The resultant images simultaneously reveal that not
only the 12 adatoms but also the 6 rest atoms per (7 × 7) unit cell of
Si(111) surface have high contrast. A careful
preparation of the STM tips (reducing the radius of the apex) may
be the key to the success, as our first-principle
calculations reveal a geometric hindrance effect of the tip apex for imaging of
such complex surfaces.


### 1.1. Experimental

The experiments were performed by using an
ultra-high-vacuum (UHV) STM system with a base
pressure of~5×10−11 mbar. The sample was an
antimony-doped n-type Si(111) wafer (resistance ~ 0.03 Ω⋅cm, thickness ~0.5 mm). Before
being introduced into the vacuum chamber, the sample is cleaned by ethanol in
an ultrasonic bath and rinsed thoroughly by deionized water. It was degassed at
about 600C° in the chamber for several
hours. Then, the sample was annealed by direct current heating while keeping
the pressure below 10−10 mbar. The annealing cycle
consisted of flashing the sample to 1200C° for 20 seconds, rapidly
lowering the temperature to about 900C°, and then slowly decreasing
the temperature at a pace of  1~2C°/s to room temperature.
Nearly perfect (7 × 7)
reconstruction was obtained by this method. Sharp STM tips made of a
polycrystalline tungsten wire were etched electrochemically in NaOH solution
and subsequently cleaned in ethanol and distilled water. Out of many tips used,
however, only three had the ability to produce repeatedly the eighteen-spot STM
images (Figure 1) while the others produce only the standard twelve-spot
images.Schematic diagram for Si(111)-7× 7 “DAS” model
[2]. (a) Top view: atoms on (111) layers with the decreasing heights indicated
by dots of decreasing sizes. The sites of corner adatom, center adatom, and
rest atom in FHUC (labeled CoA, CeA, and R, resp.), and in UHUC (labeled CoA',
CeA' and R', resp.), are identified by arrows. Positions B2 in FHUC, and B2′ in UHUC are also denoted by
arrows. (b) Side view: dangling bonds are located
at the topmost of all adatoms, rest atoms, and holes.
(a)(b)


### 1.2. Imaging Simultaneously the Rest Atoms and Adatoms

The atomic arrangement of the Si(111)-7× 7 reconstructed
surface can be described by a commonly accepted dimer-adatom-stacking (DAS)
fault model [11], as schematically shown in Figure 1. This model consists of
twelve adatoms and six rest atoms, which are evenly distributed in the faulted
half unit cell (FHUC) and unfaulted half unit cell (UHUC). Each unit cell
contains nineteen dangling bonds perpendicular to the surface, twelve for the
adatoms and six for rest atoms and one for corner atom below the vacancy. The
tunneling current in STM of the Si(111)-7 × 7 surface
originates from these dangling bonds. A stacking fault exists between the
second and third atom layers in the FHUC side (the
interlayer bonding rotates 60°), which makes the FHUC more
reactive than the UHUC. The large unit cell size (2.7 nm × 2.7 nm) makes
this surface an ideal template for the growth of well-ordered nanostructures.Figures2(a) and 2(b) show STM images with high
contrast. They demonstrate simultaneously the adatoms and the rest atoms, that
is, 18 topographic maxima per (7 × 7) unit cell.
The high-resolution image (Figure 2(b)) presents more clearly all of them. In
the UHUC side, the rest atoms appear to have almost the same brightness as the
central adatoms, whereas in the FHUC, the rest atoms appear to have
considerably less brightness than the central adatoms. The line profile in
Figure 2(c) showed the positions and height differences of the six distinct
types of atoms (labeled 1 to 6) along the solid line depicted in Figure 2(b).
The ranking of height of these atoms is as follows:
1 is the highest, then 3 and 6 follow, with 2, 4, and 5 being the lowest. The
rest atom (site 2) in the FHUC side is at the same level as the rest atom (site
5) in the UHUC side, and they are both even at the same level as the central
adatom (site 4) in the UHUC half side. The high contrast between rest atoms and
adatoms is even better than the previous results obtained by using scanning
force microscopies [5–7]. Very recently, Bassi et al. reported the
extremely similar topography of this surface at −1.5 V by using Cr tip [12],
they claimed that the Cr tip reduced its convolution effects and enhanced its
resolving capability. Here, for the first time, all the rest atoms and adatoms
of the Si(111)-7 × 7 surface are
simultaneously revealed with high contrast by the conventional W tips. The
emergence of rest atoms will be further rationalized below by theoretic
analysis.Filled-state STM images of
Si(111)-7× 7 surface
reveal 12 adatoms and 6 rest atoms per (7 × 7) unit cell.
(a) The image extends over an area of 30 nm × 30 nm. The
amplificatory (7 × 7) unit cell
was indicated in the inset. (b) Amplified image with scanning area of 8 nm × 8 nm. Both
images are recorded by sample bias voltage of −1.5 V and tunneling current of
0.3 nA. (c) The line profile taken along the line in (b).
Labels “1,” “2,” and “3” denote the corner adatom,
the rest atom, the center adatom in the FHUC, and labels “4”, “5”, and “6” denote the center adatom, the rest atom, the corner adatom in the UHUC, respectively.
(a)(b)(c)Noting the defect with the missing of one corner
adatom (close to the hole at the left-upper corner of the panel) in Figure2(b), it shows no influence on its adjacent rest atom, which is still visible
and stays its normal position without any lateral distortion. So the absence of
local adatom does not affect the geometric structures of its surrounding atoms
in (7 × 7) unit cell.
This result coincides with the recent reports about the local structures of
adatom vacancies in Si(111)-7 × 7 surface [13].
There, Chen et al. conducted STM dI/dV mappings on adatom vacancies and
found that the adatom vacancies showed different local electronic structures
but no effect to the geometric or electronic structures of the nearby rest
atoms.


### 1.3. The Emergence of Rest Atoms Is Dependent on the bias Voltage

A sequential STM snapshots obtained at different
sample bias voltages, as shown in Figure3, illustrate that the emergence of
rest atoms is dependent on the sample bias voltage. At lower bias voltages of −0.5 and −0.6 V, the images (Figures 3(a) and 3(b)) only show 12 adatoms in
each (7 × 7) unit cell.
It suggests that the electronic states of adatoms are closer to Fermi level
than those of the rest atoms. The absence of the
rest atoms ascribes to the electronic states of the rest atoms
which are outside the range of the bias when the
value of sample bias keeping very low. By further decreasing the value of bias
voltage less than −0.7 V, the rest atom spots can be visible, as shown in
Figures 3(c)–3(f). It clearly reveals that the dangling bond
states of the rest atoms are located at about 0.7 eV below the EF, which is in excellent
agreement with the experimental results measured by the method of current
imaging tunneling spectroscopy (CITS). In the year 1989, Hamers et al. measured the electronic banding structure of Si(111)-7 × 7 surface by
using CITS. They provided knowledge of the dangling bonds states of the adatoms
(about 0.35 eV below the EF) and the rest atoms (about
0.8 eV below the EF) [9]. Here, the rest atoms
appear to have almost the same brightness as the central adatoms on the UHUC
when the bias voltages are less than −0.9 V (see Figures 3(e) and 3(f)).STM images of Si(111)-7× 7 surface with
different sample bias voltages: (a) −0.5, (b) −0.6, (c) −0.7, (d) −0.8, (e)
−0.9, (f) −1.0 V, respectively. The rest atoms appear when the sample
voltages are less than −0.7 V. All images are taken at tunneling current 0.4
nA in the scanning area of 5 nm × 5 nm.
(a)(b)(c)(d)(e)(f)The STM observations presented here are in sharp
contrast to previous STM studies, which in most cases showed images similar to
that in Figure3(a) with 12 protrusions in each (7 × 7) unit,
irrespective of the bias voltages (somewhere between −2 V to 2 V). A common
explanation for the absence of the rest atom spots in the images relies on the
fact that the tunneling probability depends on the thickness of the tunneling
barrier [10]. Because the tunneling current is inversely proportional to the
exponential of the thickness, the lower-electronic state
located in the valence band corresponds to the
smaller tunneling current. The rest atoms are invisible but the adatoms are
visible may be because the former has significant lower energies than the latter.
This argument, however, contradicts the theoretical prediction that the
dangling bond states of the rest atoms extended into the vacuum region like the
adatoms [14]. Also, because the rest atoms are about 4.6 Å away from the
nearest adatoms, if one has an infinitely sharp tip positioned right above the
rest atom, there is no reason to believe that the adatoms have effect to screen
the rest-atom tunneling. If the tunneling currents
from the rest atom were indeed weak, one can move the tip closer to the surface
in a constant current STM mode. Thus, this common explanation is probably
questionable.Another possible explanation concerns tip
contamination, that is, a few silicon atoms might be accidentally picked up by
the tungsten tip during the scan, resulting in a semiconductor tip instead of
the original metallic tip. Indeed, recently it has been shown that an InAs
semiconductor tip [8] could be used to enhance rest-atom visibility by
utilizing the second gap above the fundamental gap (both lie in the Brillouin
zone center) of InAs material to suppress tunneling current from the high-lying
adatom states. However, a previous study [15] also showed that the local
electronic structure of a typical metal/semiconductor interface remains
metallic until several monolayers are in the
semiconductor. Thus, this is unlikely in the present case with Si atoms
adsorption unless the thickness of the contaminant layer exceeds the effective
screening length of Si.


### 1.4. First-Principle Calculations Are in Remarkable Agreement with Experiments

It is impractical for us to experimentally determine
what might have happened to the few tips that worked so remarkably well.
Instead, we look for a plausible explanation from theory calculations. Our
collaborators carried out the calculation by using first-principle density
function theory (DFT) [16], as implemented in the VASP codes [17]. The
Vanderbilt ultrasoft pseudopotential [18] was used with a cutoff energy equal
to 170 eV and one special k-point in the Brillouin zone sum. The surface unit
cell contains a slab of six Si layers (without counting the Si adatoms) and a
vacuum layer equivalent to six Si layers. The front surface contains the (7 × 7)
reconstruction in the Takayanagi model [11], whereas the back surface is
passivated by hydrogen. Except for the bottom layer, all the Si atoms are fully
relaxed to minimize the system total energy.Apparently, the actual tip morphology is complex,
possibly with additional atoms adsorbed at the end of the apex, as shown
schematically in the inset in Figure4(f). Because only the lower semispherical
part of the tip can be in close proximity with the surface, here the tip is
replaced by a sphere of radius r. To further simplify the calculations, only the
line-scans along the diagonal of the (7 × 7) unit cell are
considered in our simulations.(a), (c) Experimental STM images with bias
voltage of −0.57 and −1.5 V, and tunneling current of 0.3 and 0.41 nA,
respectively. F and U depict the FHUC and UHUC, respectively. (b), (d)
Calculated STM images for Si(111)-7× 7 at − 0.57 and
−1.5 V, respectively. The red peaks are about 2 Å above the dark blue
borderlines. (e), (f), and (g) are the calculated height profiles along the
diagonal of the (7 × 7) unit cell
with a tip apex radius r=0.0,7.0, and 24.0 Å, respectively. (h) The
experimental profile. The inset in (f) schematically shows an STM tip with an
adsorbed cluster beneath the apex.Figure4(a) shows the STM image of the Si(111)-7 × 7 surface at a
sample bias of −0.57 V. The appearance shows a significant contrast between the
FHUC and UHUC of the (7 × 7) unit. At
this low sample bias, the electronic states of the rest atoms are outside the
range of the bias, as demonstrated in Figures 3(a) and 3(b). Thus, the STM
topography here reveals only the twelve topmost adatoms. The adatoms in the
FHUC appear noticeably brighter than those in the UHUC. In each half, the
adatoms at the corners appear also slightly brighter than those near the
center. These qualitative features are in good agreement with the calculated
real-space charge distribution at this particular bias (Figure 4(b)). Figure 4(c) shows the STM image at a sample bias −1.5 V. Images of similar quality can
be repeatedly reproduced over large area up to 30 nm × 30 nm (Figure 2(a)). We can clearly see both the adatoms and the rest atoms. On the UHUC,
they appear to have almost the same brightness as the central adatoms, whereas
on the FUHC, the rest atoms appear to have considerably less brightness than
the central adatoms. These observations are again in excellent agreement with
the calculated real-space charge distribution at the experimental bias in
Figure 4(d).Figure4(e) shows the calculated linescan at −1.5 V
with an infinitely sharp tip, that is, r=0, as has been done before in most
STM image simulations [19]. A sharp tip is also assumed in calculating the
images in Figures 4(b) and 4(d). Now, we trace this r=0 curve with a disk of radius
r, which is a two-dimensional representation of the three-dimensional sphere,
to explore geometric hindrance. It is assumed that at each tip position,
tunneling takes place at only one spot on the disk. This is reasonable in most
cases because tunneling probability diminishes exponentially with distance.
However, there are a few exceptions where the disk is nearly or equally
distanced from the r=0 curve, that is, at or near the local symmetry points.
For simplicity, however, such a tunneling-current double effect is ignored in
our simulation.Our results show that for small disk radius mimicking
adsorbed clusters, the line-scan is essentially the same as in Figure4(e).
Figure 4(f) shows the simulated result for r=7 Å. At this radius, while none
of the main surface topological features have been lost, the overall shape of
the linescan has been significantly modified, noticeably the depth of the
profile, and the size of the atoms being noticeably larger than those in Figure 4(e). Figure 4(g) shows the simulated result for r=24 Å. At this radius, the
rest atom on the FHUC has completely vanished. Even for the UHUC, the contrast
between the rest atom spots and the adatom spots has been greatly reduced.
Thus, it is clear that the attainable size of the tip apex is the crucial
factor in imaging the true charge distribution on the (7 × 7) surfaces.
Figure 4(h) shows the corresponding linescan determined by our experiment.
Despite the simplicity of the model, the calculated result for r=7 Å in
Figure 4(f) is in quantitative agreement with experimental observation. Some of
the subtle differences between Figures 4(f) and 4(h) could probably ascribe to
the tunneling-current double effect.It is now understood that STM probes the real-space
charge distribution near theEF in a rather delicate way that
may or may not reveal the unperturbed real-space charge distribution of the
surfaces. Here, For the Si(111)-7 × 7 surface, we
show the calculated and experimental voltage-dependent charge distributions of
the Si(111)-7 × 7 surface,
which reveal simultaneously both the twelve adatoms
and six rest atoms in each (7 × 7) unit cell
[20]. The emergence of rest atom is dependent on the bias voltage and the rest
atom spots can be visible at the sample bias voltages less than −0.7 V. The
first-principle electronic structure calculations also show a strong dependence
of the charge distribution on the bias voltage: twelve spots at −0.57 V for
the twelve adatoms (see Figure 4(b)), whereas eighteen spots at −1.5 V for the
twelve adatoms plus six rest atoms (see Figure 4(d)). Our results suggest that a
geometric hindrance due to the finite size of the tip apex could be the reason
for the invisibility of the rest atoms in the past experiments. This finding
should invoke significant research interest in the design and fabrication of
the STM tip and its applications in exploring more detailed information about
surface reconstructions and nanostructures.


## 1.1. Experimental

The experiments were performed by using an
ultra-high-vacuum (UHV) STM system with a base
pressure of~5×10−11 mbar. The sample was an
antimony-doped n-type Si(111) wafer (resistance ~ 0.03 Ω⋅cm, thickness ~0.5 mm). Before
being introduced into the vacuum chamber, the sample is cleaned by ethanol in
an ultrasonic bath and rinsed thoroughly by deionized water. It was degassed at
about 600C° in the chamber for several
hours. Then, the sample was annealed by direct current heating while keeping
the pressure below 10−10 mbar. The annealing cycle
consisted of flashing the sample to 1200C° for 20 seconds, rapidly
lowering the temperature to about 900C°, and then slowly decreasing
the temperature at a pace of  1~2C°/s to room temperature.
Nearly perfect (7 × 7)
reconstruction was obtained by this method. Sharp STM tips made of a
polycrystalline tungsten wire were etched electrochemically in NaOH solution
and subsequently cleaned in ethanol and distilled water. Out of many tips used,
however, only three had the ability to produce repeatedly the eighteen-spot STM
images (Figure 1) while the others produce only the standard twelve-spot
images.Schematic diagram for Si(111)-7× 7 “DAS” model
[2]. (a) Top view: atoms on (111) layers with the decreasing heights indicated
by dots of decreasing sizes. The sites of corner adatom, center adatom, and
rest atom in FHUC (labeled CoA, CeA, and R, resp.), and in UHUC (labeled CoA',
CeA' and R', resp.), are identified by arrows. Positions B2 in FHUC, and B2′ in UHUC are also denoted by
arrows. (b) Side view: dangling bonds are located
at the topmost of all adatoms, rest atoms, and holes.
(a)(b)


## 1.2. Imaging Simultaneously the Rest Atoms and Adatoms

The atomic arrangement of the Si(111)-7× 7 reconstructed
surface can be described by a commonly accepted dimer-adatom-stacking (DAS)
fault model [11], as schematically shown in Figure 1. This model consists of
twelve adatoms and six rest atoms, which are evenly distributed in the faulted
half unit cell (FHUC) and unfaulted half unit cell (UHUC). Each unit cell
contains nineteen dangling bonds perpendicular to the surface, twelve for the
adatoms and six for rest atoms and one for corner atom below the vacancy. The
tunneling current in STM of the Si(111)-7 × 7 surface
originates from these dangling bonds. A stacking fault exists between the
second and third atom layers in the FHUC side (the
interlayer bonding rotates 60°), which makes the FHUC more
reactive than the UHUC. The large unit cell size (2.7 nm × 2.7 nm) makes
this surface an ideal template for the growth of well-ordered nanostructures.Figures2(a) and 2(b) show STM images with high
contrast. They demonstrate simultaneously the adatoms and the rest atoms, that
is, 18 topographic maxima per (7 × 7) unit cell.
The high-resolution image (Figure 2(b)) presents more clearly all of them. In
the UHUC side, the rest atoms appear to have almost the same brightness as the
central adatoms, whereas in the FHUC, the rest atoms appear to have
considerably less brightness than the central adatoms. The line profile in
Figure 2(c) showed the positions and height differences of the six distinct
types of atoms (labeled 1 to 6) along the solid line depicted in Figure 2(b).
The ranking of height of these atoms is as follows:
1 is the highest, then 3 and 6 follow, with 2, 4, and 5 being the lowest. The
rest atom (site 2) in the FHUC side is at the same level as the rest atom (site
5) in the UHUC side, and they are both even at the same level as the central
adatom (site 4) in the UHUC half side. The high contrast between rest atoms and
adatoms is even better than the previous results obtained by using scanning
force microscopies [5–7]. Very recently, Bassi et al. reported the
extremely similar topography of this surface at −1.5 V by using Cr tip [12],
they claimed that the Cr tip reduced its convolution effects and enhanced its
resolving capability. Here, for the first time, all the rest atoms and adatoms
of the Si(111)-7 × 7 surface are
simultaneously revealed with high contrast by the conventional W tips. The
emergence of rest atoms will be further rationalized below by theoretic
analysis.Filled-state STM images of
Si(111)-7× 7 surface
reveal 12 adatoms and 6 rest atoms per (7 × 7) unit cell.
(a) The image extends over an area of 30 nm × 30 nm. The
amplificatory (7 × 7) unit cell
was indicated in the inset. (b) Amplified image with scanning area of 8 nm × 8 nm. Both
images are recorded by sample bias voltage of −1.5 V and tunneling current of
0.3 nA. (c) The line profile taken along the line in (b).
Labels “1,” “2,” and “3” denote the corner adatom,
the rest atom, the center adatom in the FHUC, and labels “4”, “5”, and “6” denote the center adatom, the rest atom, the corner adatom in the UHUC, respectively.
(a)(b)(c)Noting the defect with the missing of one corner
adatom (close to the hole at the left-upper corner of the panel) in Figure2(b), it shows no influence on its adjacent rest atom, which is still visible
and stays its normal position without any lateral distortion. So the absence of
local adatom does not affect the geometric structures of its surrounding atoms
in (7 × 7) unit cell.
This result coincides with the recent reports about the local structures of
adatom vacancies in Si(111)-7 × 7 surface [13].
There, Chen et al. conducted STM dI/dV mappings on adatom vacancies and
found that the adatom vacancies showed different local electronic structures
but no effect to the geometric or electronic structures of the nearby rest
atoms.


## 1.3. The Emergence of Rest Atoms Is Dependent on the bias Voltage

A sequential STM snapshots obtained at different
sample bias voltages, as shown in Figure3, illustrate that the emergence of
rest atoms is dependent on the sample bias voltage. At lower bias voltages of −0.5 and −0.6 V, the images (Figures 3(a) and 3(b)) only show 12 adatoms in
each (7 × 7) unit cell.
It suggests that the electronic states of adatoms are closer to Fermi level
than those of the rest atoms. The absence of the
rest atoms ascribes to the electronic states of the rest atoms
which are outside the range of the bias when the
value of sample bias keeping very low. By further decreasing the value of bias
voltage less than −0.7 V, the rest atom spots can be visible, as shown in
Figures 3(c)–3(f). It clearly reveals that the dangling bond
states of the rest atoms are located at about 0.7 eV below the EF, which is in excellent
agreement with the experimental results measured by the method of current
imaging tunneling spectroscopy (CITS). In the year 1989, Hamers et al. measured the electronic banding structure of Si(111)-7 × 7 surface by
using CITS. They provided knowledge of the dangling bonds states of the adatoms
(about 0.35 eV below the EF) and the rest atoms (about
0.8 eV below the EF) [9]. Here, the rest atoms
appear to have almost the same brightness as the central adatoms on the UHUC
when the bias voltages are less than −0.9 V (see Figures 3(e) and 3(f)).STM images of Si(111)-7× 7 surface with
different sample bias voltages: (a) −0.5, (b) −0.6, (c) −0.7, (d) −0.8, (e)
−0.9, (f) −1.0 V, respectively. The rest atoms appear when the sample
voltages are less than −0.7 V. All images are taken at tunneling current 0.4
nA in the scanning area of 5 nm × 5 nm.
(a)(b)(c)(d)(e)(f)The STM observations presented here are in sharp
contrast to previous STM studies, which in most cases showed images similar to
that in Figure3(a) with 12 protrusions in each (7 × 7) unit,
irrespective of the bias voltages (somewhere between −2 V to 2 V). A common
explanation for the absence of the rest atom spots in the images relies on the
fact that the tunneling probability depends on the thickness of the tunneling
barrier [10]. Because the tunneling current is inversely proportional to the
exponential of the thickness, the lower-electronic state
located in the valence band corresponds to the
smaller tunneling current. The rest atoms are invisible but the adatoms are
visible may be because the former has significant lower energies than the latter.
This argument, however, contradicts the theoretical prediction that the
dangling bond states of the rest atoms extended into the vacuum region like the
adatoms [14]. Also, because the rest atoms are about 4.6 Å away from the
nearest adatoms, if one has an infinitely sharp tip positioned right above the
rest atom, there is no reason to believe that the adatoms have effect to screen
the rest-atom tunneling. If the tunneling currents
from the rest atom were indeed weak, one can move the tip closer to the surface
in a constant current STM mode. Thus, this common explanation is probably
questionable.Another possible explanation concerns tip
contamination, that is, a few silicon atoms might be accidentally picked up by
the tungsten tip during the scan, resulting in a semiconductor tip instead of
the original metallic tip. Indeed, recently it has been shown that an InAs
semiconductor tip [8] could be used to enhance rest-atom visibility by
utilizing the second gap above the fundamental gap (both lie in the Brillouin
zone center) of InAs material to suppress tunneling current from the high-lying
adatom states. However, a previous study [15] also showed that the local
electronic structure of a typical metal/semiconductor interface remains
metallic until several monolayers are in the
semiconductor. Thus, this is unlikely in the present case with Si atoms
adsorption unless the thickness of the contaminant layer exceeds the effective
screening length of Si.


## 1.4. First-Principle Calculations Are in Remarkable Agreement with Experiments

It is impractical for us to experimentally determine
what might have happened to the few tips that worked so remarkably well.
Instead, we look for a plausible explanation from theory calculations. Our
collaborators carried out the calculation by using first-principle density
function theory (DFT) [16], as implemented in the VASP codes [17]. The
Vanderbilt ultrasoft pseudopotential [18] was used with a cutoff energy equal
to 170 eV and one special k-point in the Brillouin zone sum. The surface unit
cell contains a slab of six Si layers (without counting the Si adatoms) and a
vacuum layer equivalent to six Si layers. The front surface contains the (7 × 7)
reconstruction in the Takayanagi model [11], whereas the back surface is
passivated by hydrogen. Except for the bottom layer, all the Si atoms are fully
relaxed to minimize the system total energy.Apparently, the actual tip morphology is complex,
possibly with additional atoms adsorbed at the end of the apex, as shown
schematically in the inset in Figure4(f). Because only the lower semispherical
part of the tip can be in close proximity with the surface, here the tip is
replaced by a sphere of radius r. To further simplify the calculations, only the
line-scans along the diagonal of the (7 × 7) unit cell are
considered in our simulations.(a), (c) Experimental STM images with bias
voltage of −0.57 and −1.5 V, and tunneling current of 0.3 and 0.41 nA,
respectively. F and U depict the FHUC and UHUC, respectively. (b), (d)
Calculated STM images for Si(111)-7× 7 at − 0.57 and
−1.5 V, respectively. The red peaks are about 2 Å above the dark blue
borderlines. (e), (f), and (g) are the calculated height profiles along the
diagonal of the (7 × 7) unit cell
with a tip apex radius r=0.0,7.0, and 24.0 Å, respectively. (h) The
experimental profile. The inset in (f) schematically shows an STM tip with an
adsorbed cluster beneath the apex.Figure4(a) shows the STM image of the Si(111)-7 × 7 surface at a
sample bias of −0.57 V. The appearance shows a significant contrast between the
FHUC and UHUC of the (7 × 7) unit. At
this low sample bias, the electronic states of the rest atoms are outside the
range of the bias, as demonstrated in Figures 3(a) and 3(b). Thus, the STM
topography here reveals only the twelve topmost adatoms. The adatoms in the
FHUC appear noticeably brighter than those in the UHUC. In each half, the
adatoms at the corners appear also slightly brighter than those near the
center. These qualitative features are in good agreement with the calculated
real-space charge distribution at this particular bias (Figure 4(b)). Figure 4(c) shows the STM image at a sample bias −1.5 V. Images of similar quality can
be repeatedly reproduced over large area up to 30 nm × 30 nm (Figure 2(a)). We can clearly see both the adatoms and the rest atoms. On the UHUC,
they appear to have almost the same brightness as the central adatoms, whereas
on the FUHC, the rest atoms appear to have considerably less brightness than
the central adatoms. These observations are again in excellent agreement with
the calculated real-space charge distribution at the experimental bias in
Figure 4(d).Figure4(e) shows the calculated linescan at −1.5 V
with an infinitely sharp tip, that is, r=0, as has been done before in most
STM image simulations [19]. A sharp tip is also assumed in calculating the
images in Figures 4(b) and 4(d). Now, we trace this r=0 curve with a disk of radius
r, which is a two-dimensional representation of the three-dimensional sphere,
to explore geometric hindrance. It is assumed that at each tip position,
tunneling takes place at only one spot on the disk. This is reasonable in most
cases because tunneling probability diminishes exponentially with distance.
However, there are a few exceptions where the disk is nearly or equally
distanced from the r=0 curve, that is, at or near the local symmetry points.
For simplicity, however, such a tunneling-current double effect is ignored in
our simulation.Our results show that for small disk radius mimicking
adsorbed clusters, the line-scan is essentially the same as in Figure4(e).
Figure 4(f) shows the simulated result for r=7 Å. At this radius, while none
of the main surface topological features have been lost, the overall shape of
the linescan has been significantly modified, noticeably the depth of the
profile, and the size of the atoms being noticeably larger than those in Figure 4(e). Figure 4(g) shows the simulated result for r=24 Å. At this radius, the
rest atom on the FHUC has completely vanished. Even for the UHUC, the contrast
between the rest atom spots and the adatom spots has been greatly reduced.
Thus, it is clear that the attainable size of the tip apex is the crucial
factor in imaging the true charge distribution on the (7 × 7) surfaces.
Figure 4(h) shows the corresponding linescan determined by our experiment.
Despite the simplicity of the model, the calculated result for r=7 Å in
Figure 4(f) is in quantitative agreement with experimental observation. Some of
the subtle differences between Figures 4(f) and 4(h) could probably ascribe to
the tunneling-current double effect.It is now understood that STM probes the real-space
charge distribution near theEF in a rather delicate way that
may or may not reveal the unperturbed real-space charge distribution of the
surfaces. Here, For the Si(111)-7 × 7 surface, we
show the calculated and experimental voltage-dependent charge distributions of
the Si(111)-7 × 7 surface,
which reveal simultaneously both the twelve adatoms
and six rest atoms in each (7 × 7) unit cell
[20]. The emergence of rest atom is dependent on the bias voltage and the rest
atom spots can be visible at the sample bias voltages less than −0.7 V. The
first-principle electronic structure calculations also show a strong dependence
of the charge distribution on the bias voltage: twelve spots at −0.57 V for
the twelve adatoms (see Figure 4(b)), whereas eighteen spots at −1.5 V for the
twelve adatoms plus six rest atoms (see Figure 4(d)). Our results suggest that a
geometric hindrance due to the finite size of the tip apex could be the reason
for the invisibility of the rest atoms in the past experiments. This finding
should invoke significant research interest in the design and fabrication of
the STM tip and its applications in exploring more detailed information about
surface reconstructions and nanostructures.


## 2. Ge Nanostructures ON Si(111)-7×7 Surfaces

Low-dimensional structures can provide interesting
physical and chemical properties due to their tiny size and shape. The growth
of nanostructures with reduced dimensions has been extensively studied, driven
by the intrinsic interest in structures as well as the potential technological
applications in quantum devices [21]. Recent studies demonstrated the
feasibilities and possibilities of growing self-organized nanostructures on
periodic solid surfaces. The Si(111)-7 × 7 surface
offers unique template for the self-assembly growth of divers nanostructures
because of the large number of distinct bonding sites. Recently, “magic”
islands and nanoclusters of semiconductor or metal have been grown on this
surface [22–26]. Ordered arrays of two-dimensional nanodots/nanoclusters,
including Al, Ga, In, Tl, Si, Ge, Sn, Pb, Na, Cu, Au, Ag, were successfully
fabricated [27–51]. These self-organized structures are expected to have a
smaller size and stronger confinement potentials compared to the
lithographically defined clusters [52].The adsorption of Germanium on the Si(111)-7× 7 surface has
been extensively studied in recent years [40–51, 53, 54, 55, 56, 57, 58, 59, 60, 61, 62, 63, 64, 65, 66, 67, 68, 69, 70, 71, 72, 73, 74, 75, 76, 77, 78, 79, 80, 81, 82, 83], because Ge-based
nanostructures have potential applications in microelectronics and
optoelectronics. Indeed, Ge/Si system naturally has
advantage compatible with Si technology. In addition, being
currently incorporated in Si structures, Ge can be used to fabricate strained
Si layers with enhanced mobility. Therefore, there are renewed activities in
Ge-based nanostructures grown on Si surface in expectation of functional
devices with unique electronic and optoelectronic properties [53].The microscopic understanding of the bonding nature of
the adsorbed Ge atoms is an essential issue for the controlled fabrication of
desired nanostructures, since the initial adsorption nature may affect the
growth behaviors of Ge-based quantum dots and films. In spite of numerous
investigations, a unified picture for the bonding structures of Ge atoms on
Si(111)-7× 7 surface has
not been established. Meanwhile, the formation and transformation process of
various Ge nanostructures during the initial growth stages is far from
being well understood. Without doubt, they impede
the further control of the growth process of Ge nanostructures.Here, we provide an STM investigation on various Ge
nanostructures on Si(111)-7× 7 surface with
different size and geometry, ranging from individual Ge atoms (adsorption
sites) to Ge nanoclusters (evolution and aggregation patterns), and then to 2D
extended Ge islands (components and bondings). Especially, we go
inside the structural characterizations as well as
the transformation process and possible mechanisms of the observed Ge nanostructures
in association with first-principle calculations.The preparation of Si(111)-7× 7 surface was
conducted as described in the part 1. Then Germanium (99.9999% purity) was
deposited onto the as-prepared Si(111)-7 × 7 surface by
resistive evaporation. The substrate was kept above room temperature (ranging
from 100 to 300C°) to facilitate the formation
of ordered structures since at room temperature, Ge atoms do not have enough
mobility to span the dimer wall after arriving on the Si(111)-7 × 7 surface [59, 61, 68]. During evaporation, the system pressure was better than  5×10−10 mbar. A typical deposition
rate of 0.01 ML/min was routinely achieved. One monolayer is defined as the
atomic density of the unreconstructed Si(111) surface (1 ML =7.83×1014 atoms/cm2). Each sample was cooled down
to room temperature, and then transferred to STM chamber for measurements. All
images were acquired in a constant-current mode with an electrochemically
etched tungsten tip.


### 2.1. Direct STM Observations of the Adsorption Sites of Ge Atoms

For the adsorption sites of Ge atoms on Si(111)-7× 7 surfaces
reported in the literatures, X-ray standing wave (XSW) studies of submonolayer
Ge deposited on Si(111)-7 × 7 at 300C° done by Patel et al. in 1985 suggested that Ge atoms might occupy substitutional-like sites on the
Si(111) plane [69]. However, the precise Ge sites and the bonding structures were
not possible to determine in their studies. Also based on XSW measurements, Dev et al. in 1986 proposed that at low coverages (~ 0.5 ML) Ge
atoms would prefer to occupy the ontop sites and to bond directly to the Si
adatoms and rest atoms which were just below the adsorbed Ge atoms [70].
Reflection electron microscopy and transmission electron diffraction
investigations on Ge/Si(111)-7 × 7 prepared at
640C° by Kajiyama et al. in
1989 found evidence that Ge atoms randomly substituted any Si atoms at the top
three layers [71]. Core-level photoemission spectroscopy measurements by
Carlisle et al. in 1994 provided indirect observations that there was
some preference for Ge to replace the Si adatoms for the annealed Ge/Si(111)-7 × 7 samples [72].
More recent measurements using near-edge X-ray absorption spectroscopy and STM
did not provide conclusive descriptions of Ge bonding sites on this surface
[56, 73, 74].Some theoretical calculations have also been reported
on Ge bonding sites on the Si(111)-7× 7 surface,
however, the calculations provided limited information and showed contradictory
results. Early work was semiempirical Xα and extended Hückel
calculations with limited predictive capabilities, which provided support for
the notion that Ge atoms bond directly to rest atoms or Si adatoms [75–77]. In
contrast to the semiempirical calculations, on the other hand, using
first-principle density functional calculations, Cho and Kaxiras in 1998
reported a limited exploration of bonding possibilities and found that the most
stable adsorption position for Ge on Si(111) is the high-coordination bridge B2 site (see Figure 1 for
pertinent terminology of the Si(111)-7 × 7 surface),
which was a bonding site that had not been proposed on the basis of
experimental data [78]. They introduced the so-called basins of attraction,
which contain stable adsorbate positions as high-coordination sites rather than
surface dangling bond sites. Their calculations showed that the rest atoms or
intrinsic Si adatoms sites (dangling bond T1 sites) of substrate were the
high-energy sites, and the low-energy sites were the B2-type sites for Si and Ge
adsorption.Here, we report STM observations and first-principle
calculations for the structure of the Ge-adsorbed Si(111)-7× 7 surface at
low Ge coverages. Figure 5 shows STM image of the Si(111)-7 × 7 surface with
Ge coverages of 0.02, 0.08, and 0.10 ML, respectively. These images show that
the surface lattice retains the original (7 × 7)
reconstruction with the dimers and the adatoms. The FHUC and the UHUC of the (7 × 7)
reconstruction are distinguished due to the different contrast. The deposited
Ge atoms appear as bright protrusions. Three significant features are presented
in the STM images. First, the deposited Ge atoms are clearly resolved as
individual atoms on the surface. Second, the adsorbed Ge atoms reside on the
sites that were occupied by the Si adatoms on Si(111)-7 × 7. Finally,
more Ge atoms occupy the corner Si adatom sites in the FHUC than the other Si
adatom sites. No Ge atoms are found at either the rest atom or the
high-coordination surface sites. Furthermore, profile lines through the bright
dots show that the height difference between the Ge atoms and the original Si
adatoms is about 0.2 Å in the STM images. This data clearly show that the Si
adatom does not stay in its original position (on a clean Si(111)-7 × 7 surface, the
Si adatom occupies a so-called T4 site just above a second-layer
Si atom) [11, 79, 80]. As the bond length of Si-Ge is about 2.36 Å, the
increased height due to addition of one Ge atom should be reflected in the STM
image.Figure 5
Filled-state STM images (20 nm× 20 nm) of the
Si(111)-7 × 7 surface with
Ge coverages of (a) 0.02 ML; (b) 0.08 ML; and (c) 0.10 ML. Sample bias: −2.2 V
in (a), and −1.5 V in (b) and (c); tunneling current: 0.5 nA in (a), and 0.2 nA
in (b) and (c). Three different configurations of Ge protrusion distributions
are denoted in (b) and (c) by solid-line triangle, dotted-line triangle, and
dashed-line triangle, respectively. The schematics for the three typical Ge
patterns, named type-A, type-B, and type-C, are shown in (d), (e), and (f),
respectively.Therefore, the addition mechanism of Ge atop Si adatom
is supposed as a questionable explanation. Moreover, the number of dangling
bonds will increase to three if a Ge atom adds on the top of one Si adatom, and
it is not considered having a suitable total energy. The topographic height
undulations of adatom sites in STM images caused by Ge-Si exchange on Si(111)-5× 5-Ge
reconstructions have been proposed by Becker et al. [81] and Fukuda [82], and were also investigated by Rosei et al. [83] with
current imaging tunneling spectroscopy. The feature of Ge-Si exchange is
confirmed by our recent results, which will be introduced in the following
section. Here, we suggest that Ge-Si exchange can also occur during the initial
adsorption stage of Ge/Si epitaxy growth due to the structures similarity of Ge
and Si. We thus conclude that Ge would prefer to substitute the Si adatoms in
its initial adsorption stages.As shown in Figures5(b) and 5(c), there are three
types of Ge protrusions patterns on the Si(111)-7 × 7 surface. The
schematics of these three types patterns, named as type-A, type-B, type-C, are
given in Figures 5(d), 5(e), and 5(f), respectively. Type-A illustrates three
Ge atoms (red spheres) locating at one corner adatom site and two adjacent
center adatom sites in a HUC. Type-B indicates the configuration with three Ge
atoms occupying corner adatom sites in a HUC. Type-C refers to the adsorption
structure with five Ge atoms residing on the sites of three corner adatoms and
two center adatoms in a HUC. Type-B and Type-C distribute preferentially in the
FHUCs, as shown in Figures 5(b) and 5(c).The sites distribution of the bright protrusions at
the corner and center adatom sites in both the FHUCs and the UHUCs is
illustrated in Figure6. At the Ge coverage of 0.02 ML, the site preference
ratio is about 5.6 : 4.4 for the FHUC to the UHUC, and 6.1 : 3.9 for the corner
to the center adatom sites, respectively. When the Ge coverage increases to
0.08 ML, the site preference ratios are about 9 : 1 for the FHUC to the UHUC, and
4 : 1 for the corner to the center adatom sites. The site distribution for the
coverage of 0.10 ML is similar to that for the coverage of 0.08 ML. The overall
conclusion is that after an initial random occupation of Si
adatom sites, corner adatom sites in the FHUC are preferred and
gradually type-B patterns become dominant. Type-A and Type-C patterns are more
discernible at slightly higher coverages, and finally, small islands begin to
appear.Figure 6
Site distributions of Ge at various adatom
positions at coverages of 0.02 ML, 0.08 ML, and 0.10 ML.Our collaborators performed first-principle DFT
calculations using the pseudopotential method and a plane-wave basis set [16, 17]. The Si(111) surface was modeled by repeated slabs with 4 layers of Si
atoms (each layer contained 16 Si atoms, corresponding to a 4 × 4 surface unit
cell) and 4 Si adatoms, separated by a vacuum region of 12 Å. Two of the four
rest atoms were saturated by hydrogen, so that the ratio of the number of the
adatoms to that of the rest atoms is the same as for the 7 × 7 surface.
Except for the Si atoms in the bottom layer, which were fixed and saturated by
H atoms, all the atoms were relaxed until the forces on them were less than
0.05 eV/Å. The exchange-correlation effects were treated with the generalized
gradient-corrected exchange-correlation functions given by Perdew and Wang
[84]. The Vanderbilt ultrasoft pseudopotentials are adopted [18]. A plane-wave
energy cutoff of 14.7 Ry and the point for reciprocal space sampling were used
for all the calculations.All the possible configurations with a Ge atom near an
adatom or/and a rest atom were calculated. Two lowest energy configurations, as
shown in Figure7, were found to have essentially the same total energy (the
difference in total energy is smaller than 0.02 eV). The first one consists of
Ge at a B2 site (Figure 7(a)), as
identified earlier by Cho and Kaxiras [85]. In the second configuration (Figure 7(b)), the adsorbed Ge atom substitutes for an Si adatom and the substituted Si
adatom occupies a nearby B2 site. We refer to the Ge
position in the second configuration as S4 (substitutional site with four
nearest-neighboring silicon atoms). The total energies of the configurations
with Ge bonded at the ontop positions of adatoms and rest atoms are significantly
higher (2.3 and 1.6 eV, resp.) than the B2 and S4 configurations. So we can
clearly rule out the possibility of such configurations, which were suggested
previously on the basis of semiempirical calculations [70, 75, 76].Schematic top view of the calculated lowest
energy configurations of a Ge atom on the Si(111) surface: (a) Ge at a B2 site and the nearby Si adatom
at the position of its original site. (b) Ge at the
substitutional S4 site and the Si adatom at a B2 site. (c) Ge atoms substitute
for some of the Si adatoms and no atoms are bonded at any of the B2 sites. The bond lengths are
shown with unit of Å.
(a)(b)(c)For both lowest energy configurations (B2 and S4), the atom (Si or Ge) at a
bridge site may diffuse within a basin (to occupy any of the six B2 sites near the rest atom) and
across basins (to occupy the B2 sites near different rest
atoms). The diffusion barriers within a basin and across basins are about 0.5 eV (0.6 eV) and 1.0 eV (1.0 eV) for the Ge (Si) atoms, respectively, which is
in agreement with previous first-principle calculations [77, 86]. Therefore, Ge
atoms in the S4 configurations are
thermodynamically more stable than in the B2 configurations. In particular,
after the atoms initially bonded at the B2 sites migrate to step edges
and/or to form islands, the surface exhibits a stable Ge-S4 configuration, in which Ge
atoms substitute for some of the Si adatoms and no atoms are bonded at any of
the B2 sites (Figure 7(c)), as shown
by our STM observations. The Ge-S4 configuration is coincided
with the recent results reported by two research groups [46, 61]. They prepared
the sample with the same experimental conditions as the current work. In their
STM measurements, they also confirmed that Ge replaced Si adatoms on the
Si(111)-7 × 7.It is well known that the backbonds of the Si adatoms
of the Si(111)-7× 7 surface are
under considerable strain [11, 79, 86]. We therefore suggest that the adsorbed
Ge atoms are able to break the backbonds and further replace the Si adatoms at
elevated temperatures. Previous studies have established that the corner
adatoms in the FHUCs are under more strain than the other adatoms, which
implies that backbonds of the corner adatoms in the FHUCs are broken easier
than those of the other adatoms [85, 86]. When Ge atoms are deposited on the
surface, the chance for the Ge atoms occupying the B2 sites near a center adatom is
larger than that near a corner adatom (the center adatom has two nearby rest
atoms while the corner adatom has only one). Thus, the Ge–S4 bonding structure tends to be
preferentially formed at the corner adatom sites and in the FHUCs of the
Si(111)-7 × 7 surface [87].Finally, the relaxed Ge–S4 configuration obtained from
our calculations shows that the Ge atom resides at the position higher by 0.24 Å than the original Si adatom that has been replaced by Ge, which is in good
agreement with our STM data.


### 2.2. Formation and Transformation of Ge Clusters

It is of course interesting to study the possible
configurations of Ge nanostructures in subsequent Ge depositions. Indeed, with
increasing Ge coverage, some novel structures, like small Ge clusters with
varying geometrical configurations appear on the Si(111)-7× 7 surface. The
representative image is shown in Figure 8(a) with Ge coverage about 0.12 ML
deposited at the substrate temperature ~ 100C°. A remarkable feature in the
image is the emergence of Ge clusters with special configurations. Dimer rows
and corner-holes of the surface are left uncovered, indicating a strong
preference of the Ge clusters to locate in (7 × 7) unit cell.
Deposited Ge clusters are imaged as bright bumps and four typical bump
structures are distinguished, as named type-Tr, type-Te, type-P, and type-H.
Figures 8(b)–8(e) show their magnified images.Figure 8
(a) Empty-state STM image (Ub = 2.0 V, It = 0.3 nA, 30 nm × 30 nm) of the
Si(111)-7 × 7 surface with
Ge coverage ~ 0.12 ML. The
substrate temperature is held at ~100C° for the Ge deposition. Ge
clusters with four typical geometrical configurations, named as type-Tr
(triangle), type-Te (tetragonal), type-P (pentagonal), and type-H (hexagonal),
are denoted by the triangles with the dotted dash-line, dotted-line,
dashed-line, and solid-line, and their magnified images (3 nm × 3 nm) are shown
in (b), (c), (d), and (e), respectively.Type-Tr (triangle-star-like) cluster emerges like a
bright triangle star in a HUC, in this structure three center adatoms and their
adjacent high-coordination sites are covered by Ge atoms. Similarly, in the
type-Te (tetragonal-star-like), type-P (pentagonal-star-like), and type-H
(hexagonal-star-like) clusters, Ge atoms occupy an increasing area in a HUC:
type-Te covers one more rest-atom region, type-P covers another one, and type-P
covers all three rest-atoms regions.Four kinds of Ge cluster configurations appear in the
same image, it suggests a formation process of Ge clusters from triangle to
tetragonal, then to pentagonal, and at last to mature hexagonal-star-like
structure despite the fact that we do not observe
the real-time evolution of single Ge cluster from simple to complex. Each kind
of Ge cluster is observed both in the FHUCs and UHUCs, as shown in Figure8(a),
and there is no clear preference for locating positions of Ge clusters in the
FHUCs and UHUCs (37 clusters in the FHUCs and 34 clusters in the UHUCs).Although the Ge cluster
structures on the Si(111)-7× 7 surface
represent the majority in Figure 8(a), while lots of individual Ge atoms
locating on the Si adatoms still can be resolved with bright spots at the
positions of some Si adatoms in the empty-state STM image. As we know, the STM
empty-state images are taken at positive sample bias voltage, corresponding to
tunneling electrons from the occupied state of tip
to sample. On clean Si(111)-7 × 7 surface, each
adatom has a dangling bond and has the same probability to accept the tunneling
electrons from tip, so all the adatoms on clean Si(111)-7 × 7 surface have
the same brightness in STM empty-state images. By this rule, we can affirm that
the brighter protrusions at the sites of Si adatoms are Ge atoms. Thus,
individual Ge atoms and some Ge clusters coexist on the Si(111)-7 × 7 surface at
proper substrate temperature and Ge coverage.


### 2.3. Evolution of Hexagonal Ge Cluster Superlattice

The above results showed that most Ge atoms form
correlated patterns at very low Ge coverages less than 0.1 ML, by replacing the
Si adatoms of the Si(111)-7× 7. With the
further increasing of Ge coverage, deposited Ge atoms are constrained inside
the HUC and aggregate into the form of clusters with different geometry and
atom numbers. In addition, previous results reported the existence of hexagonal
Ge nanostructures on Si(111)-7 × 7 [57, 60],
however, the detail process and driving force are still not clear. Here, we
will further reveal the evolution of hexagonal Ge clusters with increasing Ge
coverages. The driving mechanism and the atomic geometry of Ge clusters will be
enunciated by STM observations and first-principle calculations.In Figures9(a)–9(e), a sequence of STM images at varying
Ge coverages show the continuing evolution of the Ge clusters from isolated
ones into hexagonal patterns. The appearance illustrates that the density of Ge
clusters gradually increases with the increasing Ge coverages. The cluster
distribution becomes much regular from disordered arrangement to high-symmetric
hexagonal superlattice. Six distinct local cluster patterns can be
distinguished in STM images, as marked by different symbols. The schematics in
Figures 9(f)–9(k) simply depict the structures of these local cluster patterns,
ranging from single Ge clusters, to pair of clusters, to open cluster ring, and
finally a close cluster ring with six clusters surrounding a hole in Si(111)-7 × 7 surface. The
histograms in Figure 9(l) reveal the distribution feature of six different
local Ge nanostructures at varying coverages. The evident tendency is that the
ratio of simple cluster pattern reduces with the increase
of complex ones.Figure 9
(a)–(e) Series of STM images of Ge-deposited
Si(111)-7× 7 surface show
the formation process of hexagonal superlattice with increasing Ge coverages
ranging from 0.15, 0.2, 0.3, 0.4 to 0.5 ML, respectively. The substrate
temperature is held at ~150C° for Ge deposition. All image
sizes are 50 nm × 50 nm. The
inset in (e) shows a Fourier transform of the hexagonal arrays. (f)–(k)
Schematics illustrating the evolution of cluster structures from open to close
hexagonal ring. The Si-center adatoms that transfer charge are shaded in gray.
(l) Histograms for the distributions of different local Ge nanostructures at
varying coverages. Six distinct local nanostructures are depicted by symbols
with different shapes and colors in the STM images.The distribution tendency suggests the evolution of
hierarchical cluster patterns from dispersing clusters to close cluster rings.
Most of clusters discretely emerge on the substrate at low coverage of 0.10~ 0.15 ML with an
initial preference in the FHUCs (see Figure 9(a)). And then open cluster rings
containing three, four, and five clusters, nucleated on the Si(111)-7 × 7 surface at a
Ge coverage of ~ 0.2 ML (Figure 9(b)). Afterwards closed Ge hexagonal rings consisting of six clusters begin to
form at a Ge coverage of ~ 0.3 ML (Figure 9(c)). When the Ge coverage approaches to 0.4 ML, most of the HUCs of both FHUC
and UHUC are occupied by Ge clusters. The underlying (7 × 7) surface
periodicity and the hexagonal superstructures coexist. Finally, the
high-regular hexagonal superlattice forms at a Ge coverage of ~0.5 ML and
covers the entire (7 × 7) surface
(Figure 9(e)). The driving mechanism for the cluster evolution, ascribing to
the charge transfer from Si center adatoms to Ge clusters, will be further
discussed as following in association with DFT calculations.In Figures10(a) and 10(b), we show contrasted
filled-state STM image at sample bias −2.5 V, and empty-state image at +2.5
V, respectively. The Ge clusters look more compact in the former than in the
latter, and they show strong brightness in the center region of the FHUC. An
obvious feature is that Ge cluster has a strong effect on its three neighboring
UHUCs. The closest Si-center adatoms in the nearest-neighbor UHUCs are
invisible in the filled-state image. However, these Si adatoms do exist at
their original places in the (7 × 7)
reconstruction, as shown by the empty-state image. This fact suggests that the
Si-center adatoms in the nearest-neighbor UHUCs transfer charge to the Ge
clusters.Figure 10
Local STM images show the different
appearance of Ge-deposited Si(111)-7× 7 surface with
filled-state images (−2.5 V) in (a) and (c) and empty-state image (+2.5 V) in (b) and (d). Single Ge cluster presents in (a) and (b) and double clusters in (c) and (d). The
centers Si adatoms, as indicated by the arrows, are invisible in (a) and (c) but
visible in (b) and (d). (e) Local density of states projected onto a center Si
adatom in a UHUC before and after Ge deposition. The Fermi level is at 0 eV. (f)
The corresponding relaxed minimum energy configuration (only the FHUC is
shown). The Si and Ge atoms are depicted by gray and dark spheres,
respectively. Spheres of decreasing size represent the Si atoms with increasing
distances from the surface. The dotted lines show weak bonds.As we know, the tunneling electrons transfer from
sample to tip for the filled-state imaging, in reverse, they transfer from tip
to sample for the empty-state imaging. Here in our measurements, in the
filled-state images, the darkened areas (center adatoms) surround the clusters.
One reason is due to geometric defect, that is, no adatoms exist at these
positions. But the empty-state images prove the existence of the adatoms at the
original positions. So, the only reason is that the charge of center adatoms
transfers to nearby Ge clusters and resulting in the absence of tunneling
electrons from center adatoms in the filled-state images. Thus, the STM measurements
demonstrate lateral charge redistributions in Ge-Si system.The charge transferring from center adatoms to Ge
cluster is further revealed by first-principle DFT calculations. For all the
minimum-energy configurations, Ge clusters contain 6~ 12 Ge atoms in
the FHUC of a (7 × 7) unit cell,
and the dangling bond state of the center Si adatoms nearby the FHUC is almost
empty, indicating charge transfer of neighboring Si adatoms. Figure 10(e) shows
the projected electronic density of states (local density of states, DOS) onto
the center Si adatom in an adjacent UHUC before and after the formation of a
nine-Ge cluster in an FHUC, see the minimum-energy configuration in Figure 10(f). For clean Si(111)-7 × 7 surface
before Ge deposition, the dangling bond state of the center Si adatom is
partially occupied and crosses the Fermi level [80, 88, 89]. After the
formation of Ge clusters, the occupation of the dangling-bond state is reduced
significantly, confirming a charge transfer from the central Si adatom. Such a
charge transfer occurs because it can lower the total energy of the system.
Self-assembled clusters of various metals formed on Si(111)-7 × 7, which have
similar network as Ge clusters, have been reported [24, 27, 28, 90–92]. Several
groups suggested that the interaction between the substrate and the metal
clusters might play a role for the self-organization [24, 27, 28], and other
researchers had emphasized the interaction between the clusters themselves
[93]. Charge transfer and its role have not been reported before.When local coverage is higher, a second Ge clusters
may form at the center of an UHUC adjacent to an FHUC already containing a Ge
cluster, as shown in Figures10(c)  and 10(d). Similar to the first one, the
second Ge cluster darkens the center Si atoms in
the two neighboring FHUCs in the filled-state image (Figure 10(c)), which again
indicate charge transfer, though the charge transfer is not as effective as in
the UHUC. Charge transfer helps us understand the formation of a cluster in an
UHUC nearby an existing particle in a FHUC, and further explain the evolution
of Ge cluster patterns from isolated one to closed hexagonal patterns [94].The above results show that a new cluster forms
adjacent to an existing FHUC cluster in a neighboring UHUC, and electron
transfer occurs from the two new surrounding FHUCs. The two clusters remain
distinct with a “dimer wall” separating them as shown in Figure9(h).
First-principle calculations confirm the depletion of charge associating with
the dangling bonds of center Si adatoms in the three UHUCs surrounding a Ge
cluster in an FHUC, and also a decrease in electronic energy by such a charge
transfer. The energy gain can be attributed to “local Madelung energy”, which
is used in determining the energy of a single ion in a crystal.Assuming that the amount of charge transferring from
any of the center Si adatom is the sameq0, then a single cluster has a central charge of −3q0 (Figure 9(f)).
The total energy is lowered by a local Madelung energy of the cluster, that is
roughly (−9/d1+3/d2)q02, where d1 (~11 Å) is the
distance between the Ge clusters and an adjacent Si adatom that has been
depleted of charge, and d2 (~19 Å) is the
distance between two such Si adatoms. Because d1 is smaller than d2, the local Madelung energy is negative (−0.66q02). When a
second cluster forms in an adjacent UHUC, as in Figure 9(h), the charge on each
cluster is reduced from −3q0 to −2q0 and the local
Madelung energy is approximately  (−8/d1+2/d2+4/d3)q02≈−0.36q02, where d3 (~15.5 Å) is the
distance between the two Ge clusters. The local Madelung energy in the Ge
cluster pair is smaller than that of single cluster. The reduction in the
Madelung energy is used to overcome the factors that inhibit the formation of
isolated clusters in UHUCs, and the residual Madelung energy stabilizes the
cluster pair. The energy is substantial, and from the DOS curves of Figure 10(e), we estimate q0≈ (0.3–0.5) e, whereby the Madelung energy stabilizing a pair is ~ 0.5–1.3 eV.
With the emergence of new Ge clusters, the net
charge on each Ge cluster decreases gradually from 3q0 in the case of
an isolated cluster to 2q0, 5/3q0, 3/2q0, 7/5q0, and finally 1q0 if a complete
isolated hexagon is formed as in the schematic of Figures 9(f)–9(k), and the
effective Madelung energy per cluster also gradually reduces. So the reduction
of the local Madelung energies contributes to the stabilization of the local
cluster structures. Thus here we quantitatively revealed how the charge
transfer sustains the evolution of cluster patterns from isolated ones to
ordered hexagonal arrays.


### 2.4. Formation of Ge Islands and Ge-Si Intermixing at High Temperature

The further increase of the substrate temperature
causes the coarsening of clusters and intermixing between Ge and Si atoms,
which is believed to be due to the enhancing mobility of Ge atoms [43, 44].
When the substrate temperature is increased to about 300C° or even higher temperature,
the Ge islands begin epitaxial growth. Deposition of Ge atoms on Si(111)-7 × 7 at room
temperatures following by annealing treatment also results in 2D extended Ge
islands. Figure 11(a) shows a typical surface morphology of a Ge island with
submonolayer Ge coverage. There are three distinct features in this image.
First, the reconstruction of the island is (7 × 7), same as the
configuration of original substrate, see the close-up STM image in Figure 11(c). Second, the dimer directions of the Ge island are same as that of the
substrate which, revealing the supercell of Ge island, has the same alignment
as the substrate. Third, the shape of Ge islands is usually close to triangle,
similar to the shape of HUC triangle of the substrate. All these features
reveal the modulation effect of the substrate to the epitaxy growth of Ge
islands.(a) STM image shows a typical Ge island on
the Si(111)-7× 7 surface. The
substrate temperature was kept at 300C° for Ge deposition. (b) The
schematic drawing of the Ge island on (7 × 7)
reconstruction. (c) and (d) are amplified images of the area in (a) depicted
with a dotted-line square. (7 × 7) and (5 × 5)
reconstructions coexist in the island. (e) and (f) are the close-up images of
area in (a) depicted with a solid-line square. These images, with the irregular
distribution of the brighter atoms, illustrate the intermixing between Ge and
Si atoms. Scanning parameters: (a) 120 nm × 100 nm, 1.8 V,
0.15 nA; (c) 45 nm × 45 nm, 1.2 V,
0.15 nA; (d) 45 nm × 45 nm, −1.2 V,
0.15 nA; (e) 9 nm × 9 nm, −1.0 V,
0.15 nA; (f) 9 nm × 9 nm, −1.5 V,
0.15 nA.
(a)(b)(c)(d)(e)(f)During the growth of an island, the substrate (7× 7)
reconstruction has to be removed and the surface Si atoms will rearrange to the
bulk (1 × 1) structure
[95]. It needs to overcome different energy barriers for the removal of the
reconstruction in the UHUCs and in the FHUCs [96]. In the UHUC triangles, only
the atoms in the topmost layer rearrange, which is associated with a relatively
low-energy barrier. However, in the FHUC triangles, the removal of the stacking
faults (see the schematic in Figure 1) in the deeper layer below the adatoms is
associated with a larger energy barrier. The activation barrier for Ge
overgrowth in the FHUCs is clearly higher than in the UHUCs. Thus, the edges of
Ge island are comprised of UHUC triangles and surrounded by FHUC triangles of
the substrate, as shown in Figures 11(c) and 11(d).Figure11(b) is the schematic drawing for the Ge
island on (7 × 7)
reconstruction. FHUC triangles surrounding the Ge island as a high-energy
barrier hinder the further growth in them. Ge will nucleate on the UHUC
triangle near the FHUC triangle, as denoted by the black arrow in Figure 11(b),
where it has low energy barrier. The energy barrier
of FHUC triangle will be reduced by a gain of edge energy, thus the FHUC
triangle between the island edge and the UHUC triangle with Ge nucleation will
be attached by epitaxy Ge atoms. As a result, the Ge island shows a lateral
growth model along its edge, and the shape of Ge islands usually is triangle,
which is due to the modulation by the substrate reconstruction.Figures11(c) and 11(d) show the island involving in
several domains with two different reconstructions, (7 × 7) and (5 × 5). The domain
boundaries (defect area) are very clear. Their formation is due to the strains
between the substrate and the Ge epitaxy island. The defects on the substrate
(like the missing of adatoms and vacancies) will deform the period of (7 × 7)
reconstruction and give rise to strain [47, 97]. In addition, the mismatch of
the lattice constant of Ge and Si (Ge is 4% larger than Si) will also bring
strain. The strains can be effectively released by the formation of the domain
boundaries and the different kinds of reconstruction such as the (7 × 7) and (5 × 5) domains
shown here.The (5× 5)
reconstruction also can be described by DAS model [53, 68], as the model shown
in Figure 1 for the (7 × 7). Each (5 × 5) unit cell
includes one triangle FHUC and one UHUC, and there are three adatoms and one
rest atom distributing on the topmost layer in each HUC.The close-up filled-state STM images in Figures11(e)
and 11(f) show an irregular distribution of brighter adatoms in (7 × 7) unit cell.
The arrangement of adatoms, however, is very regular on pure Si(111)-7 × 7 surface,
where the adatoms in the FHUC are imaged brighter than the adatoms in the UHUC,
and the corner adatoms are brighter than the center adatoms in both FHUC
triangle and UHUC triangle [98]. Here, the brightness and contrast features
between adatoms disappear. In Figures 11(e) and 11(f), the corner (or center)
adatoms in the same HUC show different brightness, and even some spots at the
center adatoms sites are bright close to that of the spots at the corner
adatoms sites, so it clearly suggests the mixing condition of Ge and Si atoms.
According to the contrast feature of single Ge atoms on Si surface at very low
coverage (Figure 5), the brighter protrusions at the center adatoms sites in
Figure 11(e) are Ge atoms, and the dimmer ones are Si atoms. These observations
are coincided with the findings reported in the earlier literatures [81, 82, 99, 100], where the Ge-Si exchange in Si(111)-5 × 5 Ge
reconstructions has been proposed. Most recent
results by Voigtländer et al. provided evidences for the exchange and
intermixing of Ge/Si in Si(111)-7 × 7 surface at
high temperatures by their special techniques [43, 44]. They showed the
chemical contrast images between Si and Ge in their STM observations (Ge is
much brighter than Si) obtained on Bi-covered Ge/Si(111) surfaces. Thus, the
Ge-Si exchanging and intermixing happen at high temperature, and play an
important role in the epitaxial growth of Ge islands.In our high-temperature deposition experiments, when
Ge coverage keeps at the range of 0.2 to 0.5 ML, a novel local reconstruction
with an ordered arrangement of Ge atoms on the Si(111) surfaces is obtained.
Figure12 shows the STM images of such Ge-induced reconstruction, which
coexists with the Si(111)-7 × 7
reconstruction. The local reconstruction emerges not only inside the Ge island
(Figure 12(a)) but also inside the original (7 × 7) surface
(Figure 12(b)). The triangle domain runs over 30 nm in edge length. The
close-up image in Figure 12(c)  illustrates the local atomic structure with a
hexagonal arrangement. The atomic density is higher than the normal (7 × 7)
reconstruction and the orientation of atom rows is different from the
surrounding (7 × 7) lattice
alignment. The distance between the neighbor atoms is 0.65 ± 0.01 nm, that
is, about 3 times the
length of the basis vector (0.38 nm) for the ideal bulk-terminated Si(111)-1 × 1 unit cell. In
addition, we measured the angle between the main direction of the new local
reconstruction and the boundary of the nearby (7 × 7) unit cells
and found it to be 30°. Thus, the local
reconstruction shows a (3×3)R30° arrangement.Figure 12
STM images of 0.45 ML Ge on the Si(111)-7× 7 surface. The
substrate temperature was held at 300C° for the Ge deposition. Local
(3×3)R30° reconstruction (marked by the
white squares) emerges inside the Ge island in (a) and the Si(111)-7 × 7 substrate in
(b). (c) High-resolution image of triangle domain. (d) Schematic of the atomic
arrangement of the (3×3)R30° domain surrounded by FHUC
triangles. (e) Schematic top and side views of the atomic arrangement for the (3×3)R30° reconstruction with the
adatoms at the T4 sites. The
images are recorded at 2.0 V, 0.10 nA in (a), and 1.4 V, 0.20 nA in (b) and
(c). Image sizes: (a) 236 nm × 236 nm, (b) 123
nm × 123 nm, and (c)
22 nm × 24 nm.The appearance of images in Figures12(a) and 12(b)
demonstrates that Ge-induced  (3×3)R30° reconstruction replaces some
of the (7 × 7) unit cells,
and it does not cover the whole surface or the whole island. The local Ge
nanostructures thus coexist with the Si(111)-(7 × 7)
reconstruction. In addition, several dimmer features at some atom positions
exist within the (3×3)R30° reconstruction (Figure 12(c)),
suggesting that Si atoms are mixed with the Ge atoms.The schematic in Figure12(d) shows the atomic
arrangement of the (3×3)R30° domain
surrounded by the FHUC triangles of (7 × 7) unit cells.
As the above-mentioned analysis for the similar structure of Ge island boundary
in Figure 11(b), removing the reconstruction of the FHUC triangles requires to
overcome a larger energy barrier [96]. The activation energy for atom
rearrangement in FHUC halves is higher than that in UHUC halves. Thus, the (3×3)R30° domain propagates energy
preferentially in UHUC triangles.We further go inside the
bonding structure of the local(3×3)R30° arrangement with
support from the first-principle calculations. When
the top-layer atoms form a (3×3)R30° reconstruction on Si(111), the
underlying substrate changes its original (7 × 7)
reconstruction to (1 × 1) arrangement.
On an ideal unreconstructed Si(111) surface, there are two types of threefold
symmetric adsorption sites, known as T4, a filled position directly
above a second-layer Si atom and H3, a hollow site above a
fourth-layer Si atom sites [88, 90], as shown in Figure 12(e). The adsorbed
atoms at either T4 or H3 sites are bonded to three
first-layer Si atoms. When the dangling bonds of all the first-layer Si atoms
are saturated in this way, the adsorbed atoms form a (3×3)R30° reconstruction. Such a
reconstruction could also be formed when the adsorbed atoms occupy the
so-called S5   site (Figure 12(e)), in which
an adsorbed atom substitutes a second-layer Si atom while the replaced Si atom
is at the T4    site directly above S5 [16, 101, 102]. Our
collaborators have performed the first-principle DFT calculations for a (3×3)R30° reconstruction. In the case of
the Ge-S5 configuration, Ge or Si forms
an adlayer with a Ge coverage of 1/3 monolayer for each of the three bonding
configurations [103].The calculations show that the T4 configuration is the most
stable structure. Its total energy is lower than both the H3 and the S5 configurations by 0.60 and
0.68 eV per unit cell, respectively. This is consistent with the general
picture that the adatoms prefer to occupy the T4 sites on almost all of the
Si(111)-(3×3)R30° surfaces induced by
chemisorptions of groups III, IV, and V atoms [101]. The occurrence of Ge atoms
in the subsurface substitutional S5 sites is usually
adopted by small atoms such as boron and carbon [104–107], and is energetically
unfavorable. Occupation of a Ge atom at the subsurface S5  site would introduce
significant strain energy due to its larger size than Si. In addition, for the
Ge-S5 configuration, fully
filled Ge-associated bands do not warrant a charge transfer from the Si
dandling bond to the subsurface to decrease the surface energy as observed in
the boron-induced S5 configuration [108].
Therefore, Ge atoms would prefer to stay on the surface. While the underlying
substrate supporting the Ge-induced (3×3)R30° structure has an
unreconstructed Si(111) configuration, and significant structural relaxation is
also found.


## 2.1. Direct STM Observations of the Adsorption Sites of Ge Atoms

For the adsorption sites of Ge atoms on Si(111)-7× 7 surfaces
reported in the literatures, X-ray standing wave (XSW) studies of submonolayer
Ge deposited on Si(111)-7 × 7 at 300C° done by Patel et al. in 1985 suggested that Ge atoms might occupy substitutional-like sites on the
Si(111) plane [69]. However, the precise Ge sites and the bonding structures were
not possible to determine in their studies. Also based on XSW measurements, Dev et al. in 1986 proposed that at low coverages (~ 0.5 ML) Ge
atoms would prefer to occupy the ontop sites and to bond directly to the Si
adatoms and rest atoms which were just below the adsorbed Ge atoms [70].
Reflection electron microscopy and transmission electron diffraction
investigations on Ge/Si(111)-7 × 7 prepared at
640C° by Kajiyama et al. in
1989 found evidence that Ge atoms randomly substituted any Si atoms at the top
three layers [71]. Core-level photoemission spectroscopy measurements by
Carlisle et al. in 1994 provided indirect observations that there was
some preference for Ge to replace the Si adatoms for the annealed Ge/Si(111)-7 × 7 samples [72].
More recent measurements using near-edge X-ray absorption spectroscopy and STM
did not provide conclusive descriptions of Ge bonding sites on this surface
[56, 73, 74].Some theoretical calculations have also been reported
on Ge bonding sites on the Si(111)-7× 7 surface,
however, the calculations provided limited information and showed contradictory
results. Early work was semiempirical Xα and extended Hückel
calculations with limited predictive capabilities, which provided support for
the notion that Ge atoms bond directly to rest atoms or Si adatoms [75–77]. In
contrast to the semiempirical calculations, on the other hand, using
first-principle density functional calculations, Cho and Kaxiras in 1998
reported a limited exploration of bonding possibilities and found that the most
stable adsorption position for Ge on Si(111) is the high-coordination bridge B2 site (see Figure 1 for
pertinent terminology of the Si(111)-7 × 7 surface),
which was a bonding site that had not been proposed on the basis of
experimental data [78]. They introduced the so-called basins of attraction,
which contain stable adsorbate positions as high-coordination sites rather than
surface dangling bond sites. Their calculations showed that the rest atoms or
intrinsic Si adatoms sites (dangling bond T1 sites) of substrate were the
high-energy sites, and the low-energy sites were the B2-type sites for Si and Ge
adsorption.Here, we report STM observations and first-principle
calculations for the structure of the Ge-adsorbed Si(111)-7× 7 surface at
low Ge coverages. Figure 5 shows STM image of the Si(111)-7 × 7 surface with
Ge coverages of 0.02, 0.08, and 0.10 ML, respectively. These images show that
the surface lattice retains the original (7 × 7)
reconstruction with the dimers and the adatoms. The FHUC and the UHUC of the (7 × 7)
reconstruction are distinguished due to the different contrast. The deposited
Ge atoms appear as bright protrusions. Three significant features are presented
in the STM images. First, the deposited Ge atoms are clearly resolved as
individual atoms on the surface. Second, the adsorbed Ge atoms reside on the
sites that were occupied by the Si adatoms on Si(111)-7 × 7. Finally,
more Ge atoms occupy the corner Si adatom sites in the FHUC than the other Si
adatom sites. No Ge atoms are found at either the rest atom or the
high-coordination surface sites. Furthermore, profile lines through the bright
dots show that the height difference between the Ge atoms and the original Si
adatoms is about 0.2 Å in the STM images. This data clearly show that the Si
adatom does not stay in its original position (on a clean Si(111)-7 × 7 surface, the
Si adatom occupies a so-called T4 site just above a second-layer
Si atom) [11, 79, 80]. As the bond length of Si-Ge is about 2.36 Å, the
increased height due to addition of one Ge atom should be reflected in the STM
image.Figure 5
Filled-state STM images (20 nm× 20 nm) of the
Si(111)-7 × 7 surface with
Ge coverages of (a) 0.02 ML; (b) 0.08 ML; and (c) 0.10 ML. Sample bias: −2.2 V
in (a), and −1.5 V in (b) and (c); tunneling current: 0.5 nA in (a), and 0.2 nA
in (b) and (c). Three different configurations of Ge protrusion distributions
are denoted in (b) and (c) by solid-line triangle, dotted-line triangle, and
dashed-line triangle, respectively. The schematics for the three typical Ge
patterns, named type-A, type-B, and type-C, are shown in (d), (e), and (f),
respectively.Therefore, the addition mechanism of Ge atop Si adatom
is supposed as a questionable explanation. Moreover, the number of dangling
bonds will increase to three if a Ge atom adds on the top of one Si adatom, and
it is not considered having a suitable total energy. The topographic height
undulations of adatom sites in STM images caused by Ge-Si exchange on Si(111)-5× 5-Ge
reconstructions have been proposed by Becker et al. [81] and Fukuda [82], and were also investigated by Rosei et al. [83] with
current imaging tunneling spectroscopy. The feature of Ge-Si exchange is
confirmed by our recent results, which will be introduced in the following
section. Here, we suggest that Ge-Si exchange can also occur during the initial
adsorption stage of Ge/Si epitaxy growth due to the structures similarity of Ge
and Si. We thus conclude that Ge would prefer to substitute the Si adatoms in
its initial adsorption stages.As shown in Figures5(b) and 5(c), there are three
types of Ge protrusions patterns on the Si(111)-7 × 7 surface. The
schematics of these three types patterns, named as type-A, type-B, type-C, are
given in Figures 5(d), 5(e), and 5(f), respectively. Type-A illustrates three
Ge atoms (red spheres) locating at one corner adatom site and two adjacent
center adatom sites in a HUC. Type-B indicates the configuration with three Ge
atoms occupying corner adatom sites in a HUC. Type-C refers to the adsorption
structure with five Ge atoms residing on the sites of three corner adatoms and
two center adatoms in a HUC. Type-B and Type-C distribute preferentially in the
FHUCs, as shown in Figures 5(b) and 5(c).The sites distribution of the bright protrusions at
the corner and center adatom sites in both the FHUCs and the UHUCs is
illustrated in Figure6. At the Ge coverage of 0.02 ML, the site preference
ratio is about 5.6 : 4.4 for the FHUC to the UHUC, and 6.1 : 3.9 for the corner
to the center adatom sites, respectively. When the Ge coverage increases to
0.08 ML, the site preference ratios are about 9 : 1 for the FHUC to the UHUC, and
4 : 1 for the corner to the center adatom sites. The site distribution for the
coverage of 0.10 ML is similar to that for the coverage of 0.08 ML. The overall
conclusion is that after an initial random occupation of Si
adatom sites, corner adatom sites in the FHUC are preferred and
gradually type-B patterns become dominant. Type-A and Type-C patterns are more
discernible at slightly higher coverages, and finally, small islands begin to
appear.Figure 6
Site distributions of Ge at various adatom
positions at coverages of 0.02 ML, 0.08 ML, and 0.10 ML.Our collaborators performed first-principle DFT
calculations using the pseudopotential method and a plane-wave basis set [16, 17]. The Si(111) surface was modeled by repeated slabs with 4 layers of Si
atoms (each layer contained 16 Si atoms, corresponding to a 4 × 4 surface unit
cell) and 4 Si adatoms, separated by a vacuum region of 12 Å. Two of the four
rest atoms were saturated by hydrogen, so that the ratio of the number of the
adatoms to that of the rest atoms is the same as for the 7 × 7 surface.
Except for the Si atoms in the bottom layer, which were fixed and saturated by
H atoms, all the atoms were relaxed until the forces on them were less than
0.05 eV/Å. The exchange-correlation effects were treated with the generalized
gradient-corrected exchange-correlation functions given by Perdew and Wang
[84]. The Vanderbilt ultrasoft pseudopotentials are adopted [18]. A plane-wave
energy cutoff of 14.7 Ry and the point for reciprocal space sampling were used
for all the calculations.All the possible configurations with a Ge atom near an
adatom or/and a rest atom were calculated. Two lowest energy configurations, as
shown in Figure7, were found to have essentially the same total energy (the
difference in total energy is smaller than 0.02 eV). The first one consists of
Ge at a B2 site (Figure 7(a)), as
identified earlier by Cho and Kaxiras [85]. In the second configuration (Figure 7(b)), the adsorbed Ge atom substitutes for an Si adatom and the substituted Si
adatom occupies a nearby B2 site. We refer to the Ge
position in the second configuration as S4 (substitutional site with four
nearest-neighboring silicon atoms). The total energies of the configurations
with Ge bonded at the ontop positions of adatoms and rest atoms are significantly
higher (2.3 and 1.6 eV, resp.) than the B2 and S4 configurations. So we can
clearly rule out the possibility of such configurations, which were suggested
previously on the basis of semiempirical calculations [70, 75, 76].Schematic top view of the calculated lowest
energy configurations of a Ge atom on the Si(111) surface: (a) Ge at a B2 site and the nearby Si adatom
at the position of its original site. (b) Ge at the
substitutional S4 site and the Si adatom at a B2 site. (c) Ge atoms substitute
for some of the Si adatoms and no atoms are bonded at any of the B2 sites. The bond lengths are
shown with unit of Å.
(a)(b)(c)For both lowest energy configurations (B2 and S4), the atom (Si or Ge) at a
bridge site may diffuse within a basin (to occupy any of the six B2 sites near the rest atom) and
across basins (to occupy the B2 sites near different rest
atoms). The diffusion barriers within a basin and across basins are about 0.5 eV (0.6 eV) and 1.0 eV (1.0 eV) for the Ge (Si) atoms, respectively, which is
in agreement with previous first-principle calculations [77, 86]. Therefore, Ge
atoms in the S4 configurations are
thermodynamically more stable than in the B2 configurations. In particular,
after the atoms initially bonded at the B2 sites migrate to step edges
and/or to form islands, the surface exhibits a stable Ge-S4 configuration, in which Ge
atoms substitute for some of the Si adatoms and no atoms are bonded at any of
the B2 sites (Figure 7(c)), as shown
by our STM observations. The Ge-S4 configuration is coincided
with the recent results reported by two research groups [46, 61]. They prepared
the sample with the same experimental conditions as the current work. In their
STM measurements, they also confirmed that Ge replaced Si adatoms on the
Si(111)-7 × 7.It is well known that the backbonds of the Si adatoms
of the Si(111)-7× 7 surface are
under considerable strain [11, 79, 86]. We therefore suggest that the adsorbed
Ge atoms are able to break the backbonds and further replace the Si adatoms at
elevated temperatures. Previous studies have established that the corner
adatoms in the FHUCs are under more strain than the other adatoms, which
implies that backbonds of the corner adatoms in the FHUCs are broken easier
than those of the other adatoms [85, 86]. When Ge atoms are deposited on the
surface, the chance for the Ge atoms occupying the B2 sites near a center adatom is
larger than that near a corner adatom (the center adatom has two nearby rest
atoms while the corner adatom has only one). Thus, the Ge–S4 bonding structure tends to be
preferentially formed at the corner adatom sites and in the FHUCs of the
Si(111)-7 × 7 surface [87].Finally, the relaxed Ge–S4 configuration obtained from
our calculations shows that the Ge atom resides at the position higher by 0.24 Å than the original Si adatom that has been replaced by Ge, which is in good
agreement with our STM data.


## 2.2. Formation and Transformation of Ge Clusters

It is of course interesting to study the possible
configurations of Ge nanostructures in subsequent Ge depositions. Indeed, with
increasing Ge coverage, some novel structures, like small Ge clusters with
varying geometrical configurations appear on the Si(111)-7× 7 surface. The
representative image is shown in Figure 8(a) with Ge coverage about 0.12 ML
deposited at the substrate temperature ~ 100C°. A remarkable feature in the
image is the emergence of Ge clusters with special configurations. Dimer rows
and corner-holes of the surface are left uncovered, indicating a strong
preference of the Ge clusters to locate in (7 × 7) unit cell.
Deposited Ge clusters are imaged as bright bumps and four typical bump
structures are distinguished, as named type-Tr, type-Te, type-P, and type-H.
Figures 8(b)–8(e) show their magnified images.Figure 8
(a) Empty-state STM image (Ub = 2.0 V, It = 0.3 nA, 30 nm × 30 nm) of the
Si(111)-7 × 7 surface with
Ge coverage ~ 0.12 ML. The
substrate temperature is held at ~100C° for the Ge deposition. Ge
clusters with four typical geometrical configurations, named as type-Tr
(triangle), type-Te (tetragonal), type-P (pentagonal), and type-H (hexagonal),
are denoted by the triangles with the dotted dash-line, dotted-line,
dashed-line, and solid-line, and their magnified images (3 nm × 3 nm) are shown
in (b), (c), (d), and (e), respectively.Type-Tr (triangle-star-like) cluster emerges like a
bright triangle star in a HUC, in this structure three center adatoms and their
adjacent high-coordination sites are covered by Ge atoms. Similarly, in the
type-Te (tetragonal-star-like), type-P (pentagonal-star-like), and type-H
(hexagonal-star-like) clusters, Ge atoms occupy an increasing area in a HUC:
type-Te covers one more rest-atom region, type-P covers another one, and type-P
covers all three rest-atoms regions.Four kinds of Ge cluster configurations appear in the
same image, it suggests a formation process of Ge clusters from triangle to
tetragonal, then to pentagonal, and at last to mature hexagonal-star-like
structure despite the fact that we do not observe
the real-time evolution of single Ge cluster from simple to complex. Each kind
of Ge cluster is observed both in the FHUCs and UHUCs, as shown in Figure8(a),
and there is no clear preference for locating positions of Ge clusters in the
FHUCs and UHUCs (37 clusters in the FHUCs and 34 clusters in the UHUCs).Although the Ge cluster
structures on the Si(111)-7× 7 surface
represent the majority in Figure 8(a), while lots of individual Ge atoms
locating on the Si adatoms still can be resolved with bright spots at the
positions of some Si adatoms in the empty-state STM image. As we know, the STM
empty-state images are taken at positive sample bias voltage, corresponding to
tunneling electrons from the occupied state of tip
to sample. On clean Si(111)-7 × 7 surface, each
adatom has a dangling bond and has the same probability to accept the tunneling
electrons from tip, so all the adatoms on clean Si(111)-7 × 7 surface have
the same brightness in STM empty-state images. By this rule, we can affirm that
the brighter protrusions at the sites of Si adatoms are Ge atoms. Thus,
individual Ge atoms and some Ge clusters coexist on the Si(111)-7 × 7 surface at
proper substrate temperature and Ge coverage.


## 2.3. Evolution of Hexagonal Ge Cluster Superlattice

The above results showed that most Ge atoms form
correlated patterns at very low Ge coverages less than 0.1 ML, by replacing the
Si adatoms of the Si(111)-7× 7. With the
further increasing of Ge coverage, deposited Ge atoms are constrained inside
the HUC and aggregate into the form of clusters with different geometry and
atom numbers. In addition, previous results reported the existence of hexagonal
Ge nanostructures on Si(111)-7 × 7 [57, 60],
however, the detail process and driving force are still not clear. Here, we
will further reveal the evolution of hexagonal Ge clusters with increasing Ge
coverages. The driving mechanism and the atomic geometry of Ge clusters will be
enunciated by STM observations and first-principle calculations.In Figures9(a)–9(e), a sequence of STM images at varying
Ge coverages show the continuing evolution of the Ge clusters from isolated
ones into hexagonal patterns. The appearance illustrates that the density of Ge
clusters gradually increases with the increasing Ge coverages. The cluster
distribution becomes much regular from disordered arrangement to high-symmetric
hexagonal superlattice. Six distinct local cluster patterns can be
distinguished in STM images, as marked by different symbols. The schematics in
Figures 9(f)–9(k) simply depict the structures of these local cluster patterns,
ranging from single Ge clusters, to pair of clusters, to open cluster ring, and
finally a close cluster ring with six clusters surrounding a hole in Si(111)-7 × 7 surface. The
histograms in Figure 9(l) reveal the distribution feature of six different
local Ge nanostructures at varying coverages. The evident tendency is that the
ratio of simple cluster pattern reduces with the increase
of complex ones.Figure 9
(a)–(e) Series of STM images of Ge-deposited
Si(111)-7× 7 surface show
the formation process of hexagonal superlattice with increasing Ge coverages
ranging from 0.15, 0.2, 0.3, 0.4 to 0.5 ML, respectively. The substrate
temperature is held at ~150C° for Ge deposition. All image
sizes are 50 nm × 50 nm. The
inset in (e) shows a Fourier transform of the hexagonal arrays. (f)–(k)
Schematics illustrating the evolution of cluster structures from open to close
hexagonal ring. The Si-center adatoms that transfer charge are shaded in gray.
(l) Histograms for the distributions of different local Ge nanostructures at
varying coverages. Six distinct local nanostructures are depicted by symbols
with different shapes and colors in the STM images.The distribution tendency suggests the evolution of
hierarchical cluster patterns from dispersing clusters to close cluster rings.
Most of clusters discretely emerge on the substrate at low coverage of 0.10~ 0.15 ML with an
initial preference in the FHUCs (see Figure 9(a)). And then open cluster rings
containing three, four, and five clusters, nucleated on the Si(111)-7 × 7 surface at a
Ge coverage of ~ 0.2 ML (Figure 9(b)). Afterwards closed Ge hexagonal rings consisting of six clusters begin to
form at a Ge coverage of ~ 0.3 ML (Figure 9(c)). When the Ge coverage approaches to 0.4 ML, most of the HUCs of both FHUC
and UHUC are occupied by Ge clusters. The underlying (7 × 7) surface
periodicity and the hexagonal superstructures coexist. Finally, the
high-regular hexagonal superlattice forms at a Ge coverage of ~0.5 ML and
covers the entire (7 × 7) surface
(Figure 9(e)). The driving mechanism for the cluster evolution, ascribing to
the charge transfer from Si center adatoms to Ge clusters, will be further
discussed as following in association with DFT calculations.In Figures10(a) and 10(b), we show contrasted
filled-state STM image at sample bias −2.5 V, and empty-state image at +2.5
V, respectively. The Ge clusters look more compact in the former than in the
latter, and they show strong brightness in the center region of the FHUC. An
obvious feature is that Ge cluster has a strong effect on its three neighboring
UHUCs. The closest Si-center adatoms in the nearest-neighbor UHUCs are
invisible in the filled-state image. However, these Si adatoms do exist at
their original places in the (7 × 7)
reconstruction, as shown by the empty-state image. This fact suggests that the
Si-center adatoms in the nearest-neighbor UHUCs transfer charge to the Ge
clusters.Figure 10
Local STM images show the different
appearance of Ge-deposited Si(111)-7× 7 surface with
filled-state images (−2.5 V) in (a) and (c) and empty-state image (+2.5 V) in (b) and (d). Single Ge cluster presents in (a) and (b) and double clusters in (c) and (d). The
centers Si adatoms, as indicated by the arrows, are invisible in (a) and (c) but
visible in (b) and (d). (e) Local density of states projected onto a center Si
adatom in a UHUC before and after Ge deposition. The Fermi level is at 0 eV. (f)
The corresponding relaxed minimum energy configuration (only the FHUC is
shown). The Si and Ge atoms are depicted by gray and dark spheres,
respectively. Spheres of decreasing size represent the Si atoms with increasing
distances from the surface. The dotted lines show weak bonds.As we know, the tunneling electrons transfer from
sample to tip for the filled-state imaging, in reverse, they transfer from tip
to sample for the empty-state imaging. Here in our measurements, in the
filled-state images, the darkened areas (center adatoms) surround the clusters.
One reason is due to geometric defect, that is, no adatoms exist at these
positions. But the empty-state images prove the existence of the adatoms at the
original positions. So, the only reason is that the charge of center adatoms
transfers to nearby Ge clusters and resulting in the absence of tunneling
electrons from center adatoms in the filled-state images. Thus, the STM measurements
demonstrate lateral charge redistributions in Ge-Si system.The charge transferring from center adatoms to Ge
cluster is further revealed by first-principle DFT calculations. For all the
minimum-energy configurations, Ge clusters contain 6~ 12 Ge atoms in
the FHUC of a (7 × 7) unit cell,
and the dangling bond state of the center Si adatoms nearby the FHUC is almost
empty, indicating charge transfer of neighboring Si adatoms. Figure 10(e) shows
the projected electronic density of states (local density of states, DOS) onto
the center Si adatom in an adjacent UHUC before and after the formation of a
nine-Ge cluster in an FHUC, see the minimum-energy configuration in Figure 10(f). For clean Si(111)-7 × 7 surface
before Ge deposition, the dangling bond state of the center Si adatom is
partially occupied and crosses the Fermi level [80, 88, 89]. After the
formation of Ge clusters, the occupation of the dangling-bond state is reduced
significantly, confirming a charge transfer from the central Si adatom. Such a
charge transfer occurs because it can lower the total energy of the system.
Self-assembled clusters of various metals formed on Si(111)-7 × 7, which have
similar network as Ge clusters, have been reported [24, 27, 28, 90–92]. Several
groups suggested that the interaction between the substrate and the metal
clusters might play a role for the self-organization [24, 27, 28], and other
researchers had emphasized the interaction between the clusters themselves
[93]. Charge transfer and its role have not been reported before.When local coverage is higher, a second Ge clusters
may form at the center of an UHUC adjacent to an FHUC already containing a Ge
cluster, as shown in Figures10(c)  and 10(d). Similar to the first one, the
second Ge cluster darkens the center Si atoms in
the two neighboring FHUCs in the filled-state image (Figure 10(c)), which again
indicate charge transfer, though the charge transfer is not as effective as in
the UHUC. Charge transfer helps us understand the formation of a cluster in an
UHUC nearby an existing particle in a FHUC, and further explain the evolution
of Ge cluster patterns from isolated one to closed hexagonal patterns [94].The above results show that a new cluster forms
adjacent to an existing FHUC cluster in a neighboring UHUC, and electron
transfer occurs from the two new surrounding FHUCs. The two clusters remain
distinct with a “dimer wall” separating them as shown in Figure9(h).
First-principle calculations confirm the depletion of charge associating with
the dangling bonds of center Si adatoms in the three UHUCs surrounding a Ge
cluster in an FHUC, and also a decrease in electronic energy by such a charge
transfer. The energy gain can be attributed to “local Madelung energy”, which
is used in determining the energy of a single ion in a crystal.Assuming that the amount of charge transferring from
any of the center Si adatom is the sameq0, then a single cluster has a central charge of −3q0 (Figure 9(f)).
The total energy is lowered by a local Madelung energy of the cluster, that is
roughly (−9/d1+3/d2)q02, where d1 (~11 Å) is the
distance between the Ge clusters and an adjacent Si adatom that has been
depleted of charge, and d2 (~19 Å) is the
distance between two such Si adatoms. Because d1 is smaller than d2, the local Madelung energy is negative (−0.66q02). When a
second cluster forms in an adjacent UHUC, as in Figure 9(h), the charge on each
cluster is reduced from −3q0 to −2q0 and the local
Madelung energy is approximately  (−8/d1+2/d2+4/d3)q02≈−0.36q02, where d3 (~15.5 Å) is the
distance between the two Ge clusters. The local Madelung energy in the Ge
cluster pair is smaller than that of single cluster. The reduction in the
Madelung energy is used to overcome the factors that inhibit the formation of
isolated clusters in UHUCs, and the residual Madelung energy stabilizes the
cluster pair. The energy is substantial, and from the DOS curves of Figure 10(e), we estimate q0≈ (0.3–0.5) e, whereby the Madelung energy stabilizing a pair is ~ 0.5–1.3 eV.
With the emergence of new Ge clusters, the net
charge on each Ge cluster decreases gradually from 3q0 in the case of
an isolated cluster to 2q0, 5/3q0, 3/2q0, 7/5q0, and finally 1q0 if a complete
isolated hexagon is formed as in the schematic of Figures 9(f)–9(k), and the
effective Madelung energy per cluster also gradually reduces. So the reduction
of the local Madelung energies contributes to the stabilization of the local
cluster structures. Thus here we quantitatively revealed how the charge
transfer sustains the evolution of cluster patterns from isolated ones to
ordered hexagonal arrays.


## 2.4. Formation of Ge Islands and Ge-Si Intermixing at High Temperature

The further increase of the substrate temperature
causes the coarsening of clusters and intermixing between Ge and Si atoms,
which is believed to be due to the enhancing mobility of Ge atoms [43, 44].
When the substrate temperature is increased to about 300C° or even higher temperature,
the Ge islands begin epitaxial growth. Deposition of Ge atoms on Si(111)-7 × 7 at room
temperatures following by annealing treatment also results in 2D extended Ge
islands. Figure 11(a) shows a typical surface morphology of a Ge island with
submonolayer Ge coverage. There are three distinct features in this image.
First, the reconstruction of the island is (7 × 7), same as the
configuration of original substrate, see the close-up STM image in Figure 11(c). Second, the dimer directions of the Ge island are same as that of the
substrate which, revealing the supercell of Ge island, has the same alignment
as the substrate. Third, the shape of Ge islands is usually close to triangle,
similar to the shape of HUC triangle of the substrate. All these features
reveal the modulation effect of the substrate to the epitaxy growth of Ge
islands.(a) STM image shows a typical Ge island on
the Si(111)-7× 7 surface. The
substrate temperature was kept at 300C° for Ge deposition. (b) The
schematic drawing of the Ge island on (7 × 7)
reconstruction. (c) and (d) are amplified images of the area in (a) depicted
with a dotted-line square. (7 × 7) and (5 × 5)
reconstructions coexist in the island. (e) and (f) are the close-up images of
area in (a) depicted with a solid-line square. These images, with the irregular
distribution of the brighter atoms, illustrate the intermixing between Ge and
Si atoms. Scanning parameters: (a) 120 nm × 100 nm, 1.8 V,
0.15 nA; (c) 45 nm × 45 nm, 1.2 V,
0.15 nA; (d) 45 nm × 45 nm, −1.2 V,
0.15 nA; (e) 9 nm × 9 nm, −1.0 V,
0.15 nA; (f) 9 nm × 9 nm, −1.5 V,
0.15 nA.
(a)(b)(c)(d)(e)(f)During the growth of an island, the substrate (7× 7)
reconstruction has to be removed and the surface Si atoms will rearrange to the
bulk (1 × 1) structure
[95]. It needs to overcome different energy barriers for the removal of the
reconstruction in the UHUCs and in the FHUCs [96]. In the UHUC triangles, only
the atoms in the topmost layer rearrange, which is associated with a relatively
low-energy barrier. However, in the FHUC triangles, the removal of the stacking
faults (see the schematic in Figure 1) in the deeper layer below the adatoms is
associated with a larger energy barrier. The activation barrier for Ge
overgrowth in the FHUCs is clearly higher than in the UHUCs. Thus, the edges of
Ge island are comprised of UHUC triangles and surrounded by FHUC triangles of
the substrate, as shown in Figures 11(c) and 11(d).Figure11(b) is the schematic drawing for the Ge
island on (7 × 7)
reconstruction. FHUC triangles surrounding the Ge island as a high-energy
barrier hinder the further growth in them. Ge will nucleate on the UHUC
triangle near the FHUC triangle, as denoted by the black arrow in Figure 11(b),
where it has low energy barrier. The energy barrier
of FHUC triangle will be reduced by a gain of edge energy, thus the FHUC
triangle between the island edge and the UHUC triangle with Ge nucleation will
be attached by epitaxy Ge atoms. As a result, the Ge island shows a lateral
growth model along its edge, and the shape of Ge islands usually is triangle,
which is due to the modulation by the substrate reconstruction.Figures11(c) and 11(d) show the island involving in
several domains with two different reconstructions, (7 × 7) and (5 × 5). The domain
boundaries (defect area) are very clear. Their formation is due to the strains
between the substrate and the Ge epitaxy island. The defects on the substrate
(like the missing of adatoms and vacancies) will deform the period of (7 × 7)
reconstruction and give rise to strain [47, 97]. In addition, the mismatch of
the lattice constant of Ge and Si (Ge is 4% larger than Si) will also bring
strain. The strains can be effectively released by the formation of the domain
boundaries and the different kinds of reconstruction such as the (7 × 7) and (5 × 5) domains
shown here.The (5× 5)
reconstruction also can be described by DAS model [53, 68], as the model shown
in Figure 1 for the (7 × 7). Each (5 × 5) unit cell
includes one triangle FHUC and one UHUC, and there are three adatoms and one
rest atom distributing on the topmost layer in each HUC.The close-up filled-state STM images in Figures11(e)
and 11(f) show an irregular distribution of brighter adatoms in (7 × 7) unit cell.
The arrangement of adatoms, however, is very regular on pure Si(111)-7 × 7 surface,
where the adatoms in the FHUC are imaged brighter than the adatoms in the UHUC,
and the corner adatoms are brighter than the center adatoms in both FHUC
triangle and UHUC triangle [98]. Here, the brightness and contrast features
between adatoms disappear. In Figures 11(e) and 11(f), the corner (or center)
adatoms in the same HUC show different brightness, and even some spots at the
center adatoms sites are bright close to that of the spots at the corner
adatoms sites, so it clearly suggests the mixing condition of Ge and Si atoms.
According to the contrast feature of single Ge atoms on Si surface at very low
coverage (Figure 5), the brighter protrusions at the center adatoms sites in
Figure 11(e) are Ge atoms, and the dimmer ones are Si atoms. These observations
are coincided with the findings reported in the earlier literatures [81, 82, 99, 100], where the Ge-Si exchange in Si(111)-5 × 5 Ge
reconstructions has been proposed. Most recent
results by Voigtländer et al. provided evidences for the exchange and
intermixing of Ge/Si in Si(111)-7 × 7 surface at
high temperatures by their special techniques [43, 44]. They showed the
chemical contrast images between Si and Ge in their STM observations (Ge is
much brighter than Si) obtained on Bi-covered Ge/Si(111) surfaces. Thus, the
Ge-Si exchanging and intermixing happen at high temperature, and play an
important role in the epitaxial growth of Ge islands.In our high-temperature deposition experiments, when
Ge coverage keeps at the range of 0.2 to 0.5 ML, a novel local reconstruction
with an ordered arrangement of Ge atoms on the Si(111) surfaces is obtained.
Figure12 shows the STM images of such Ge-induced reconstruction, which
coexists with the Si(111)-7 × 7
reconstruction. The local reconstruction emerges not only inside the Ge island
(Figure 12(a)) but also inside the original (7 × 7) surface
(Figure 12(b)). The triangle domain runs over 30 nm in edge length. The
close-up image in Figure 12(c)  illustrates the local atomic structure with a
hexagonal arrangement. The atomic density is higher than the normal (7 × 7)
reconstruction and the orientation of atom rows is different from the
surrounding (7 × 7) lattice
alignment. The distance between the neighbor atoms is 0.65 ± 0.01 nm, that
is, about 3 times the
length of the basis vector (0.38 nm) for the ideal bulk-terminated Si(111)-1 × 1 unit cell. In
addition, we measured the angle between the main direction of the new local
reconstruction and the boundary of the nearby (7 × 7) unit cells
and found it to be 30°. Thus, the local
reconstruction shows a (3×3)R30° arrangement.Figure 12
STM images of 0.45 ML Ge on the Si(111)-7× 7 surface. The
substrate temperature was held at 300C° for the Ge deposition. Local
(3×3)R30° reconstruction (marked by the
white squares) emerges inside the Ge island in (a) and the Si(111)-7 × 7 substrate in
(b). (c) High-resolution image of triangle domain. (d) Schematic of the atomic
arrangement of the (3×3)R30° domain surrounded by FHUC
triangles. (e) Schematic top and side views of the atomic arrangement for the (3×3)R30° reconstruction with the
adatoms at the T4 sites. The
images are recorded at 2.0 V, 0.10 nA in (a), and 1.4 V, 0.20 nA in (b) and
(c). Image sizes: (a) 236 nm × 236 nm, (b) 123
nm × 123 nm, and (c)
22 nm × 24 nm.The appearance of images in Figures12(a) and 12(b)
demonstrates that Ge-induced  (3×3)R30° reconstruction replaces some
of the (7 × 7) unit cells,
and it does not cover the whole surface or the whole island. The local Ge
nanostructures thus coexist with the Si(111)-(7 × 7)
reconstruction. In addition, several dimmer features at some atom positions
exist within the (3×3)R30° reconstruction (Figure 12(c)),
suggesting that Si atoms are mixed with the Ge atoms.The schematic in Figure12(d) shows the atomic
arrangement of the (3×3)R30° domain
surrounded by the FHUC triangles of (7 × 7) unit cells.
As the above-mentioned analysis for the similar structure of Ge island boundary
in Figure 11(b), removing the reconstruction of the FHUC triangles requires to
overcome a larger energy barrier [96]. The activation energy for atom
rearrangement in FHUC halves is higher than that in UHUC halves. Thus, the (3×3)R30° domain propagates energy
preferentially in UHUC triangles.We further go inside the
bonding structure of the local(3×3)R30° arrangement with
support from the first-principle calculations. When
the top-layer atoms form a (3×3)R30° reconstruction on Si(111), the
underlying substrate changes its original (7 × 7)
reconstruction to (1 × 1) arrangement.
On an ideal unreconstructed Si(111) surface, there are two types of threefold
symmetric adsorption sites, known as T4, a filled position directly
above a second-layer Si atom and H3, a hollow site above a
fourth-layer Si atom sites [88, 90], as shown in Figure 12(e). The adsorbed
atoms at either T4 or H3 sites are bonded to three
first-layer Si atoms. When the dangling bonds of all the first-layer Si atoms
are saturated in this way, the adsorbed atoms form a (3×3)R30° reconstruction. Such a
reconstruction could also be formed when the adsorbed atoms occupy the
so-called S5   site (Figure 12(e)), in which
an adsorbed atom substitutes a second-layer Si atom while the replaced Si atom
is at the T4    site directly above S5 [16, 101, 102]. Our
collaborators have performed the first-principle DFT calculations for a (3×3)R30° reconstruction. In the case of
the Ge-S5 configuration, Ge or Si forms
an adlayer with a Ge coverage of 1/3 monolayer for each of the three bonding
configurations [103].The calculations show that the T4 configuration is the most
stable structure. Its total energy is lower than both the H3 and the S5 configurations by 0.60 and
0.68 eV per unit cell, respectively. This is consistent with the general
picture that the adatoms prefer to occupy the T4 sites on almost all of the
Si(111)-(3×3)R30° surfaces induced by
chemisorptions of groups III, IV, and V atoms [101]. The occurrence of Ge atoms
in the subsurface substitutional S5 sites is usually
adopted by small atoms such as boron and carbon [104–107], and is energetically
unfavorable. Occupation of a Ge atom at the subsurface S5  site would introduce
significant strain energy due to its larger size than Si. In addition, for the
Ge-S5 configuration, fully
filled Ge-associated bands do not warrant a charge transfer from the Si
dandling bond to the subsurface to decrease the surface energy as observed in
the boron-induced S5 configuration [108].
Therefore, Ge atoms would prefer to stay on the surface. While the underlying
substrate supporting the Ge-induced (3×3)R30° structure has an
unreconstructed Si(111) configuration, and significant structural relaxation is
also found.


## 3. Conclusions

Firstly, we reported UHV-STM experiments and
first-principle total energy calculations which are combined to determine the
STM images of Si(111)-7×7 surface. Both the rest atoms and adatoms were observed
simultaneously with high contrast by using the conventional W tips. The
emergence of the rest atoms was dependent on the sample bias voltage. The rest
atom spots could be visible at the bias voltages less than −0.7 V, and their
brightness is even comparable to that of the center Si adatoms when the voltage
is less than −0.9 V. The possible explanations for
the visibility of rest atoms in our STM images were discussed and a very
sharper tip could resolve them, which were enunciated by first-principle
calculations.Secondly, we investigated the structural
characterizations and the bonding nature of diverse Ge nanostructures on
Si(111)-7×7 surface at different deposition stages. We performed STM
measurements of the adsorption site of single Ge atom on the Si(111)-7×7 surfaces for a sequence of submonolayer coverages deposited
at 150C°. The observations suggested
that individual Ge atoms replaced the so-called Si
adatoms rather than being adsorbed directly atop of
the Si adatoms. Initially, the replacements were random, but distinct patterns
emerged when increasing the Ge coverages, until small clusters
are formed on the substrate. The first-principle
density-functional calculations revealed that Ge/Si substitution configuration
was more energetically favorable and thermodynamically stable than the
arrangements of Ge locating at the high-coordination surface sites.Further deposited Ge atoms generated nanoclusters with
varying geometrical configurations. Individual Ge atoms and Ge clusters
coexisted on the Si(111)-7×7 surfaces. Ge nanoclusters gradually produced in both the
faulted and unfaulted half unit cells of (7×7) units with an initial preference in the faulted halves,
and ultimately self-organized into the form of well-ordered hexagonal
superlattice corresponding to the geometry of one Ge cluster per triangle half
unit of original (7×7) lattice. Charge transfer from Si adatoms to Ge
nanoclusters played a key role in the self-organization of the superlattice,
which was proved by experimental observations and theoretic calculations.Two-dimensional extended Ge islands with triangle
shape were formed on the substrate when its temperature was kept at 300C° for Ge deposition. The irregular distribution of brighter
topmost adatoms suggested the intermixing status of Ge/Si components in the
islands, and the intermixing ascribed to the exchanging of Ge atoms with the
substrate Si atoms at higher temperatures. Several local domains with different
reconstructions like (5×5) and (3×3)R30°arrangements were
found on the substrates. The configuration of the Ge adatoms residing at the T4 sites rather than S5 or H3 positions in the (3×3)R30°reconstruction was
proposed according to the first-principle calculations.



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*Source: 874213-2008-04-14.xml*