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Journal of Physics: Condensed Matter

1

TOPICAL REVIEW

Growth morphology of thin films on the metallic and oxide

surfaces

Aleksander Krupski

Department of Physics, University of Warwick, Coventry CV4 7AL, United Kingdom.

E-mail: [email protected]

Abstract

In this work we briefly review recent investigations concerning growth morphology of thin metallic

films on the Mo(110) and the Ni3Al(111) surfaces, and Fe and Copper Phthalocyanine (C32H16N8Cu)

on the Al2O3/Ni3Al(111) surface. Comparison of Ag, Au, Sn, and Pb growth on the Mo(110) surface

has shown a number of similarities between these adsorption systems except surface alloy formation

that has only observed in the case of Sn and Au. In the Pb/Mo(110) and Pb/Ni3Al(111) adsorption

systems selective formation of uniform Pb island heights during metal thin film growth has been

observed and interpreted in terms of quantum size effects. Furthermore, our studies showed that Al2O3

on Ni3Al(111) exhibits a large superstructure in which the unit cell has a commensurate relation to the

substrate lattice. In addition, Copper Phthalocyanine chemisorbed weakly onto an ultrathin Al2O3 film

on Ni3Al(111) and showed a poor template effect of the Al2O3/Ni3Al(111) system. In the case of iron

cluster growth on Al2O3/Ni3Al(111) the nucleation sites were independent of deposition temperature,

yet cluster shape showed a dependence. In this system, Fe clusters formed a regular hexagonal lattice

on the Al2O3/Ni3Al(111).

PACS:

68.55.-a, 68.55.J-, 68.37.Ef, 68.47.De, 68.47.Gh, 68.47.Jn, 68.43.Bc, 68.43.Fg, 68.43.Mn, 71.20.-b,

68.60.Dv, 68.35.B-, 68.35.bd, 71.15.-m, 61.05.jh

Keywords:

Aluminum; Nickel; Aluminum oxide; Iron; Copper Phthalocyanine; Molybdenum; Lead; Silver; Gold;

Tin; STM; STS; LEED; SPA-LEED; AES; Ab initio; Density functional theory; Calculations; Growth;

Thin film growth; Cluster growth and nucleation; Metal–metal interfaces; Electron–solid interactions;

Low index single crystal surface; Surface alloys.


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Contents

1. Introduction 2

2. Thin metal films on the Mo(110) and

Ni3Al(111) surfaces 4

2.1. Mo(110) 4

2.2. Ni3Al(111) 4

2.3. Ag/Mo(110) 7

2.4. Au/Mo(110) 10

2.5. Sn/Mo(110) 11

2.6. Pb/Mo(110) 13

2.7. Pb/Ni3Al(111) 16

3. Molecules and metal growth on the

Al2O3/Ni3Al(111) 19

3.1. Al2O3/Ni3Al(111) 19

3.2. C32H16N8Cu/Al2O3/Ni3Al(111) 21

3.3. Fe clusters on Al2O3/Ni3Al(111) 24

4. Conclusions 26

Acknowledgements 27

References 27

1. Introduction

Ultra-thin epitaxial film systems exhibit a variety

of interesting properties owing to the strong

correlation between the electronic structure of the

film and its morphology, strain, and defect

structure [1-2]. Knowledge of the interface

structure between face-centered cubic fcc thin

metal films and body-centered cubic bcc metals

is important in understanding the initial growth of

these adsorption systems. Structural studies of

fcc/bcc systems provide a great deal of

information on the connection between

geometrical properties of the adsorbed atomic

layers and atomic arrangements of the substrates.

Until now, the surface structures and growth

modes of various adsorbate metals (e.g. Cu [3],

Co [4], Fe [5], Ni [6], Mg [7], Ba [8], K [9], S

[10], In [11], Ag [12-17], Au [18-19], Sn [20-21],

Pb [22–25], etc.) on the Mo(110) substrates have

been investigated by a variety of experimental

and theoretical methods in a number of works.

Recently, the formation of oxide layers on

transition metal (TM) surfaces has also received

considerable attention [26-29] including Mo(110)

surface [30-31]. In addition, Pb layers also been

of interest in recent years because the islands

formed during growth are of a special uniform

height (often referred to as a ‘magic’ or a

‘wedding cake’ island) with flat tops and sharp

step edges. This phenomenon has been observed

during metal thin-film growth in a number of

material systems such as Pb/Mo(110) [22],

Pb/Ni3Al(111) [32-33], Pb/Ni(111) [34-38],

Pb/Cu(111) [39-42], Pb/Si(111) [41,43-55],

In/Si(111)-Pb-α-3×3 [56], Ag/Fe(001) [57],

Ag/GaAs(110) [58], Ag/Si(111)-7×7 [59] and has

been interpreted in terms of quantum size effects

(QSE) [60-64]. Quantum size effects were

initially observed experimentally by Jaklevic et

al. [65-67] for metals during measurements of

tunnelling currents in metal-metal and oxide-

metal junctions and later by Jonker et al. [68-70]

in low-energy electron transmission experiments.

Since then, QSE-related oscillations have been

found in other transport property measurements

of thin films (e.g. superconducting films, the Hall

coefficient, and the electric resistivity), all of

which depend on the density of states at the

Fermi energy. QSE studies indicate that, in the

equilibrium distribution of island heights, some

heights appear much more frequently than others.

‘Magic’ preferred heights correspond to islands

with a quantum well state far from the Fermi

energy, while ‘forbidden’ heights appear to be

those that have a quantum well state close to the

Fermi level. When these uniform height islands

are observed the growth conditions for

temperature and flux are different for each

substrate type suggesting that kinetic factors play

a role. ’-Ni3Al is of great interest for its

attractive applications in the power and aerospace

industries as a high-temperature structural

material. In a bulk or thin-film state this

compound exhibits unique physical and

mechanical properties such as a high melting

point, high hardness, high yield strength, low

density and good oxidation and corrosion

resistance [71-77]. Nickel aluminate surfaces,

including the Ni3Al(111) and Ni3Al(001)

surfaces, have been investigated experimentally

[78-82] and theoretically [81-91], and are often

used as a substrate in studies of a variety of

different surface phenomena such as oxidation

(e.g. Al2O3/Ni3Al(111) [92], Al2O3/NiAl(110)

[93-94]) and the formation of more complex

systems like CuPc/Al2O3/Ni3Al(111) [95] or

Fe/Al2O3/Ni3Al(111) [96]. Ni3Al can be seen as a

model system for use as a template for a well-

ordered Al2O3 surface [92], which is suitable for

studies with several surface sensitive

investigation methods requiring electric

conductivity. Aside from the usage as a support


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material for heterogeneous catalysts [97], Al2O3

is also of increasing relevance as an insulator for

semiconductor devices because of its superior

dielectric properties compared to the classical

insulator SiO2 [98].

The investigation of organic thin films grown on

semiconducting surfaces has already been of

interest for many years. The high thermal

stability and electronic properties of

phthalocyanine (Pc) compounds and their metal

complexes (e.g. CuPc) are very interesting

candidates for the development of new technical

applications in various fields; examples include

highly specific gas sensors [99-101], organic

light emitting diodes [102], organic field effect

transistors [103-104] and photovoltaic cells

[105]. The main focus of our work with the

CuPc/Al2O3/Ni3Al(111) system [95] was to

determine whether the well-established template

effect of the Al2O3/Ni3Al(111) system for metal

cluster growth [106-107] could also be exploited

for an ordered arrangement of organic molecules.

This ordering is extremely important for the

development of novel technologies. Since the

first scanning tunneling microscopy (STM)

investigations of CuPc on a polycrystalline Ag

film by Gimzewski et al. [108] many different

substrates have been considered for the growth of

ordered CuPc films. In the initial work it was not

possible to reach submolecular resolution and no

surface induced orientation of the molecules were

achieved by Lippel et al. on Cu(100) [109].

However, the formation of an ordered CuPc

monolayer was found later on other metal

surfaces like Ag(111) [110] and Au(111) [111].

Some semiconducting substrates like graphite

and MoS2 also were observed to exhibit an

ordered planar growth of the molecules [112].

The first STM and scanning tunneling

spectroscopy (STS) investigations of CuPc on an

ultrathin alumina film on Ni3Al(110) by Ho and

co-workers [113] did not show a significant

template effect of the substrate but the authors

could distinguish between two different kinds of

adsorbate configurations with D4h- and D2h-

symmetry, respectively. Finally, it is necessary to

understand metal-oxide interfaces for progression

of many technologies including metal-ceramic

sensors, microelectronics, spintronic devices and

oxide supported transition metal catalysts [114-

119]. In our work concerning

Fe/Al2O3/Ni3Al(111) [96] we focus on the

creation of small oxide supported iron clusters

serving as a model system to investigate the

catalytic and magnetic properties of nanosized

metallic particles on an insulating substrate. The

catalytic properties of Fe clusters on oxides are of

particular interest with regard to nitrogen

fixation, being of vital importance in biological

systems and of the highest relevance in

technological processes [120]. Biological

nitrogen fixation was the subject of numerous

experimental as well as theoretical studies aimed

at creating a detailed picture of the catalytic

activity of an enzyme and of the reaction

pathway [121-123]. A recent publication

investigating small Fe clusters on MgO(100)

attributes to this model system catalytic

properties that are similar to the active sites in the

biological enzyme nitrogenase [124]. The effects

of the reduced particle size and the insulating

substrate are not only crucial for catalytic activity

but are expected to introduce new physical

properties. The magnetic properties of the small

iron particles supported by oxide systems are of

particular interest with respect to the

development of high-density magnetic data

recording devices [125-126]. Density functional

calculations revealed that small Fen clusters

(n ≤ 8) deposited on an inert substrate have

magnetic moments largely exceeding the Fe bulk

value [127], which is consistent with the same

observation for metal clusters in the gas phase

[128].

In this article we review our recent investigations

concerning growth morphology of thin metallic

films (Ag, Au, Sn) on Mo(110), Pb on Mo(110)

and Ni3Al(111) surfaces and Fe, Copper

Phthalocyanine (C32H16N8Cu) on

Al2O3/Ni3Al(111). Our studies were performed

with the use of well know and widely described

experimental methods such as: STM/STS [129-

136], low-energy electron diffraction (LEED)

[137] including spot profile analysis low-energy

electron diffraction (SPA-LEED) [138-139],

Auger electron spectroscopy (AES) [140] and

theoretical calculations based on the density

functional theory (DFT) [141] and plane-wave

basis set, as implemented in the Vienna Ab initio

Simulation Package (VASP) [142-145].

Presented STM data were processed by freeware

image-processing software [146].

The rest of this article is divided into the

following sections. Section 2 will give important

information on thin metal films on Mo(110) and

Ni3Al(111) surfaces. Section 3 molecules and

metal growth on the Al2O3/Ni3Al(111) are

discussed. The conclusion will consist of a brief

summary of the previous sections.


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2. Thin metal films on the Mo(110) and

Ni3Al(111) surfaces

2.1. Mo(110)

Mo has a body-centred cubic (bcc) crystal

structure with the space group 𝐼𝑚3̅𝑚 with lattice

parameter a = 3.15 Å [147-148], see Fig.1(a).

The structural model for the clean Mo(110)

surface and the corresponding LEED pattern are

shown in Fig. 1(b) and Fig. 1(c), respectively.

Fig.1(d) displays an STM image that was taken

on a low-index Mo(110) substrate with terraces

between 100 and 400 Å wide separated by

monoatomic steps aligned along the ]111[

direction. The height of the steps on the Mo(110)

surface was measured by STM to be 2.3 ± 0.2 Å.

Theoretical studies of the Mo(110) surface [23]

performed with the use of ab initio calculations

based on the density functional theory and plane-

wave basis set, as implemented in the VASP

package indicate that the relaxation process of the

Mo(110) surface system is limited solely to its

topmost part. The first interlayer distance d1,2 was

calculated to be reduced by 5% (i.e., somewhat

above 0.1 Å) while the subsequent distance d2,3

increased by 0.5–0.9 % (0.01–0.03 Å). Following

this, the changes in the third interlayer distance

d3,4 was seen already as negligibly small. These

results were in good agreement with earlier

theoretical studies [149–152] as well as LEED

measurements [153]. The results presented in

[23] showed that the increase in the slab

thickness above seven layers did not considerably

change the atomic structure of the surface. Local

density-of-states (LDOS) distributions that were

calculated for the three topmost atomic layers of

Mo(110) indicate that electronic states in the

energy region around the Fermi level were

distinctly localized at the surface atomic layer

while at lower energies their localization was

shifted toward subsurface layers. LDOS

calculations performed for the topmost atomic

layer shows that the surface electronic structure

of Mo(110) is dominated by d states within the

whole energy range that was investigated [23].

2.2. Ni3Al(111)

Ni3Al has an A3B-L12- or Cu3Au-type structure

with space group 𝑃𝑚3̅𝑚, and the equilibrium

lattice parameter a = 3.57 Å [154-155], see

Fig.2(a). The structural model for the clean

Ni3Al(111) surface and corresponding LEED

Figure 1. (a) Schematic view of the bulk unit cell of Mo. The (110) plane is marked. (b) Top view of Mo(110)

surface. (c) LEED pattern of the clean Mo(110) surface recorded for E = 300 eV at T = 300 K, and normal

electron incidence. The (11) unit cell of the substrate is drawn in yellow. (d) STM image of the clean Mo(110)

at T = 300 K (1000 Å × 1000Å, IT = 0.8 nA, Ubias = 100 mV). The line scan shown in the inset evidences

terraces approximately between 100 and 400 Å wide. (d) is taken from [23]. Copyright (2009) by The

American Physical Society.


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pattern are shown in Fig. 2(b) and Fig. 2(c),

respectively. The face-centred cubic (fcc) (111)

surface of this perfectly ordered alloy, shown in

Fig.2(b), has exactly the same stoichiometric

composition as the bulk. Each aluminium atom is

embedded in the nickel matrix. This results in a

2×2 aluminium arrangement in the surface layer.

The lattice constant of the surface unit cell is 5.04

Å (double the parameter of the primitive fcc

(111) unit cell). Fig. 2(d) shows an atomically

resolved STM image of the clean (111) surface of

Ni3Al. The image shows hexagonal arrangements

with a lattice constant of about 5.074 Å. This

distance corresponds well with the (2×2)

superstructure of the surface unit cell due to the

ordered distribution of Al and Ni atoms [154].

The nearest Ni-Ni distance at the surface is 2.54

Å and cannot be seen in STM images.

Nevertheless, as it was shown in work of L.

Jurczyszyn et al. [156], this fact is not

condradictory to the bulk-like stoichiometric

composition of the surface layer. However,

instead of a long-range periodicity, the STM

image showed relatively small domains with a

highly visible (2×2) order, see Fig.2(d). Our

study indicates that such a short-range order of

the surface structure is an intristic property of the

Ni3Al(111) system. The bias voltage in this

experiment was +50mV, and it is important to

stress that the experimental setup requires this

voltage to be applied to the sample. This means

that at the given voltage tunnelling occurs from

the tip into unoccupied states of the sample. It

followed from our experimental and theoretical

studies [81] that the topographies of Ni3Al(111)

STM images were strongly influenced by intra-

atomic s-pz interference. This kind of interference

increased the efficiency of electron tunneling

through s and pz orbitals of surface Al atoms,

while reducing considerably the corresponding

current contributions flowing through the surface

Ni atoms. Earlier studies [156] indicated that this

Figure 2. (a) Schematic view of the bulk unit cell of L12-ordered Ni3Al. The (111) and (001) planes are

outlined. (b) Top view of Ni3Al(111) surface. (c) LEED pattern of the clean p(2×2)-Ni3Al(111) surface

recorded for E = 136 eV at T = 200 K, and normal electron incidence. (d) STM image of the clean p(2×2)-

Ni3Al(111) at T = 23 K (75 Å × 75 Å, IT = 260 pA, Ubias = 50 mV). Tunneling spectra measured at the points

indicated on the STM image. Curves 1–5 were acquired after the feedback loop was opened at 85 mV at a

tunneling current of 120 pA. (e) STM image of the Al2O3/Ni3Al(111) at T = 23 K (71 Å × 71 Å, IT = 72 pA,

Ubias = 3.21 V). Tunneling spectra measured at the points indicated on the STM image. Curves 1–5 were

acquired after the feedback loop was opened at 190 mV at a tunneling current of 72 pA. The corresponding

modulation frequency was 5 kHz with an amplitude of 20 mV. Solid red line shows the DOS calculated for the

Ni3Al(111) surface. The (11) and (22) unit cells are drawn in blue and red, respectively. (c) is adapted from

[32]. Copyright (2013) by Elsevier. (d) and (e) are taken from [81]. Copyright (2009) by The American

Physical Society.


6

factor might have been responsible for the

domination of surface Al atoms in STM images.

In particular, it was found that even for an

unrelaxed surface structure, where surface Al and

Ni atoms have the same vertical positions (as it

takes place in the bulk), the topography of the

calculated STM image was dominated by Al

atoms. More specifically, the surface Al atoms

appeared in the simulated STM images of the

(2×2) superstructure with the same shape and

size as they appeared in experimental STM data,

while the surface Ni atoms were completely

invisible. In the experimentally measured relaxed

Ni3Al(111) surface structure, the domination of

Al atoms in STM images was additionally

supported by the differences in vertical positions

of Al and Ni atoms [81]. Fig.2(e) shows the STS

spectra obtained for five different horizontal

positions of the tip, marked by arrows in Fig.2(d).

This set of spectra indicate that, except for the

maximum located at the Fermi level, the energies

of the rest of STS features dominating in the

considered energy range somewhat depend on the

position of the STM tip. Apart from the

experimental data, Fig.2(e) also presents the

LDOS distribution at the surface layer obtained

from our present ab initio calculations (solid red

line). It should be pointed out here that in the

standard STS measurements, the obtained spectra

mainly reflected the energy distributions of s and

p states of the substrate surface, as d states are

not very active during the tunneling process.

Therefore, for the comparison with experimental

results, the LDOS distribution plotted in Fig.2(d)

takes into account only the contributions

connected with s and p states of the surface

atoms. This theoretical result represents the

averaged LDOS distribution over three Ni atoms

and one Al atom from the (2×2) surface unit cell.

Below the Fermi level, the theoretical distribution

indicates the presence of LDOS features located

0.4 eV and 0.6 eV below EF. The lower one (at

−0.6 eV) corresponds well with a distinct STS

peak that appears around −0.6 eV in each

experimental curve. However, the higher

theoretical feature (at −0.4 eV) can only be linked

with a relatively weak shoulder, which can be

seen in STS spectra 1, 2, and 4 around −0.4 eV.

Dependencies presented in Fig.2(d) also indicate

that theoretical results reproduce the measured

STS features in the vicinity of the Fermi level

well. For unoccupied states, the calculated LDOS

distribution shows a distinct feature located

around 0.35 eV. Such a feature can be seen in

STS spectra 2 and 5, though in spectrum 2 it is

somewhat shifted toward higher energies. Much

weaker maxima within this energy range can also

be observed in curves 1 and 3. However, as

mentioned earlier, the experimental studies of

Ni3Al(111) indicated a lack of the long-range

order along this surface (see Figs.2(d)) while the

theoretical study was performed assuming ideal

translation symmetry. This fact needs to be taken

into account when comparing STS spectra and

LDOS distributions presented in Fig.2(e).

Previous STM/STS study [157] showed that for

the Al2O3/Ni3Al(111) system, the ordering of the

substrate underneath the oxide film was much

higher than for the clean Ni3Al(111) surface.

Properties of the Al2O3/Ni3Al(111) system are

described in more detail in section 4.1 of this

review. The STS data indicated the presence of a

large oxide band gap [118, 157-158] that allowed

STM/STS measurements of the Ni3Al(111)

substrate surface through the oxide film. Fig. 2(e)

shows an STM image of Al2O3/Ni3Al(111)

collected with a bias of 3.2 V, i.e., when the

topography of the image was dominated by the

parameters of the Al2O3 structure. However, the

set of STS spectra presented in Fig.2(e) was

Figure 3. (a) STM image of the clean Ni3Al(001) at T = 4 K (100 Å × 100 Å, IT = 5 nA, Ubias = 85 mV). The FFT

shown in the inset evidences a square lattice arrangement of protrusions (Al atoms) with the nearest neighbor

distance of 3.7 Å. (b) Top view of Ni3Al(001) surface. (c) LEED pattern of the clean Ni3Al(001) surface recorded

for E = 120 eV, and normal electron incidence. The unit cells are outlined. (a) and (c) are adapted from [82].

Copyright (2008) by Elsevier.


7

obtained for the bias range −1.5 to 1.5 V, which

corresponds to the energy gap of the oxide. Since

for such voltages the Al2O3 film appeared

invisible in STM/STS measurements [81], a

reduction of the tip sample separation enabled us

to interrogate the electronic structure of the

Ni3Al(111) substrate. Fig.2(e) shows a set of STS

spectra measured at five different horizontal

positions of the STM tip, marked by arrows in

Fig.2(e). This sequence of dependencies shows

that, contrary to the case of a clean Ni3Al(111),

see Fig. 2(d), the STS spectra had very little

dependence on the horizontal position of the tip.

We may say that for energies within the oxide

band gap, the presence of the oxide film does not

modify the surface electronic structure of the

substrate much, as compared to the

corresponding results for a clean Ni3Al(111)

surface. In contrast to the Ni3Al(111) surface, the

Ni3Al(001) surface [82] clearly exhibits a long-

range character, see Fig.3(a). The structural

model of the clean Ni3Al(001) surface and

corresponding LEED pattern are presented in Fig.

3(b) and Fig. 2(c), respectively. Topographies

from the obtained STM images indicate a surface

superstructure with a square lattice unit and a

lattice constant equal to (= 3.7 Å) the distance

between the nearest neighbour surface atoms of

the same type (i.e., nearest Al or nearest Ni

atoms). It indicates that the STM image is formed

by one type of surface atom only, while the

atoms of the second type remain invisible. It was

shown in [143] that the STM process at the

Ni3Al(001) surface is strongly influenced by the

intra-atomic s-pz interface. This kind of

interference reduces the s and pz current

contributions flowing through Ni surface atoms

and increases considerably the s-pz contributions

connected with surface Al atoms. A detailed

analysis has shown that this factor leads to Al

atoms being the dominant species in simulated

STM images of the Ni3Al(001) surface [82,156].

2.3. Ag/Mo(110)

It is well know that plotting AES peak heights of

the substrate and of the adsorbate as a function of

deposition time (AES(t) plots) enables the

determination of the growth mechanism as well

as monolayer formation [33, 160-164]. In our

studies at room temperature, only one linear

region in the AES(t) plots for 186 eV

molybdenum and 351 eV silver peaks could be

distinguished, see Fig. 4(a). It could be deduced

from this region in the AES(t) plots that two-

dimensional growth of the first silver layer is

observed. Subsequently, the non-linear shape of

the AES(t) plots suggested the occurrence of

three-dimensional (3D) growth. If A

S = hS1/ hS0

defines the coefficient of attenuation of the

substrate Auger peak due to the presence of a

monolayer of adsorbate, then the expected height

of the substrate Auger peak for layer-by-layer

growth after completion of the n = 2nd, 3rd, 4th, ...

layer is given by the equation hSn = hS0 (A

S )n

[33]. Curve (1) in Fig. 1 was calculated for FM

growth using the formula described above, under

the assumption that at the (hS1; t1) = (0.555; 198)

point the first silver layer was completed. It can

be seen from a comparison of our calculated and

experimental AES(t) plots that two-dimensional

growth of the second silver layer should have

also been observed. However, one cannot

recognize any distinct breaking points by eye

after deposition times t2, t3 or t4 on the

experimentally found AES(t) plot. Fitting an

exponential to the data provided a closer

agreement. Such a shape of an AES(t) plot

suggests the Stranski-Krastanov growth mode

(3D). Because the scatter of hS1 and t1 values for

these AES(t) plots were not large, we will further

assume that the initial linear part of the curve

corresponds to the formation of the first layer (θ

= 1 ML) of silver, where a 1 ML of Ag(111) film

corresponds to an atomic packing density of

13.801014 atoms/cm2. Owing to the method used

for the recording of AES data, the dependencies

of AES signals for both substrate and adsorbate

on adsorbate deposition time were represented by

as many as 20 to 30 experimental data points for

each monolayer of adsorbate. Thus, profiles of

AES kinetics could be determined more precisely

than for the kinetics presented in [13] where one

monolayer was represented by only 14 points and

only by adsorbate AES kinetics. In addition, in

our measurements, adsorbate deposition was not

interrupted for the recording of Auger peaks.

Thus, our AES kinetics analysis was for

continuous growth of the deposited layer. Fig.

4(b)-(d) shows STM images of the Mo(110)

surface taken with different Ag coverages in

order to illustrate the development of

morphology of the Ag layers deposited on Mo at

room temperature. Fig. 4(b) shows a typical STM

image corresponding to a submonolayer coverage

of θ = 0.2 ML. The bright features represent one

monolayer high silver islands. The observed Ag

islands on the substrate had regular shapes, sharp

edges, and smooth tops. At this point, an analysis


8

of the STM measurements indicated that, for

coverage less than 1 ML, the growth of Ag was

mediated by a two dimensional step-flow

mechanism. This means that, during the initial

stage of growth, the first Ag film nucleates and

creates islands at Mo step edges. That is, the

growth proceeds by incorporation of Ag adatoms

at ascending step edges and the film advances

across the atomic terraces along the mis-cut

direction. The estimated average size of the

islands is about 180 20 nm2. Fig. 4(c) shows an

STM image corresponding to monolayer

coverage θ = 1.0 ML, as deduced from the

AES(t) plot in Fig. 4(a). However, the first two-

dimensional silver layer was not complete as was

expected from the AES(t) plots. At this coverage,

the two-dimensional step-flow mechanism was

continued. The silver islands of the first layer

were then larger than in the submonolayer

regime. In addition, between these silver islands,

additional small Ag islands without flat tops

appeared. Also, the second silver layer started to

grow before completion of the first silver layer,

Fig. 4(c). In this image the first and second layer

is marked by numbers I and II, respectively.

Arrows in this image point to Ag islands which

extended over two terraces of the substrate, and

their thickness did not change by 1 ML from

terrace to terrace. This island shape type has been

termed a “nano-wedge island” [165]. Closer

inspection of Fig. 4(c) shows that, in the

monolayer coverage range, the decoration of the

substrate steps by Ag could be distinguished by

the presence of a fractional step of p1 = 0.86 ± 0.6

Å height at the Ag-Mo boundary. This fractional

step may have been due to a difference in either

the local density of states (Mo(110) [166-167],

Ag [168]) or the atomic radii of the two metals

(ar, Mo = 1.40 Å, ar,Ag = 1.44 Å). At coverages

greater than θ = 7 ML, see Fig. 4(d), further

formation of three dimensional (3D) islands with

sharp edges and mostly smooth tops was

Figure 4. Ag deposition on Mo(110) face at T = 300 K. (a) AES(t) plot of Mo MNN and Ag MNN peak

heights. (1) AES(t) plot calculated for Frank–van der Merwe growth. STM images at coverages: (b) 0.20

ML (4010 Å × 4010 Å, IT = 2.85 nA, Ubias = 0.1 mV). (c) 1.00 ML (1000 Å × 1000 Å, IT = 6.97 nA, Ubias =

20 mV). Arrows in this image point to Ag islands which extend over two terraces of the substrate, and their

thickness does not change by 1 ML from terrace to terrace. (d) 7 ML (800 Å × 800 Å, IT = 0.2 nA, Ubias =

2.0 V). The arrows indicate the positions where the Mo substrate is still visible. (e) post annealed 1.0 ML

at T = 700 K (5047 Å × 5047 Å, IT = 0.2 nA, Ubias = 2.5 V). (f) Line-scan corresponding to line A drawn in (e)

demonstrating that the Ag nanostripes protrude by 0.39 0.06 Å with respect to a single step of the Mo

terrace. The insets in (d) and (e) present differentiated STM images with enhanced contrast. Adapted from

[12]. Copyright (2010) by Elsevier.


9

observed. The silver islands probably minimize

their surface free energy by having flat tops. The

arrows in Fig. 4(d) indicate the areas where the

Mo substrate was still visible. These STM results

agreed well with our AES studies and confirmed

a 3D growth mode, a result that contradicted the

literature [13-17] where layer-by-layer growth at

room temperature up to 3 ML was reported. In

addition, the results presented here are in good

agreement with a simulated growth mode [169],

where a 3D growth mode was shown with island

growth commencing from the second Ag layer. A

thermodynamic description of growth is based on

the specific surface free energy of a particular

crystal facet, defined as the reversible work per

unit area required for the creation of the surface

[170]. Different growth modes may be

distinguished according to a balance between

surface free energy of substrate (S) and film (A)

and interface energy (I). Namely, layer-by-layer

or Frank-Van der Merwe (FM) growth is

expected if S ≥ A + I. In hetero-epitaxial

systems, the surface free energy increases with

the number of film layers (n) so that the FM

condition will no longer be valid at a critical film

thickness (n*) at which point three–dimensional

(3D) crystals begin to form (Stranski-Krastanov

(SK) mode). The Volumer-Weber (VW) mode,

where 3D growth begins immediately,

corresponds to n* = 1. The higher surface energy

of the Mo substrate allows the Ag film to grow in

pseudomorphic registry with the substrate. This

was confirmed by Mae [169], using a molecular

dynamics aided kinetic Monte Carlo method.

Table 1 contains set of structural parameters and

surface energies () of Mo, Ni3Al and also Pb,

Ag, Au, Sn with respect to Miller plane

orientation reproduced from [44,147-

148,154,171-180]. When the sample was post-

annealed to 700 K, the morphology of the surface

changed, as shown in Figs. 4(e)–(f). At this

temperature, growth was driven by a pure step-

flow mechanism. Step-flow growth in the

submonolayer and monolayer regime gives rise

to a regular Ag nanostripe network attached to

Mo(110) step edges, Fig. 4(e) for 1.0 ML. The

corrugation profiles (A) in Fig. 4(f) highlight the

protrusion of silver nanostripes equal to p2 =

0.39± 0.06 Å for 1.0 ML with respect to a single

step of a Mo terrace. In addition, the width and

shape of the top of the nanostripes showed a

dependence on coverage. It is suggested that the

variations seen in morphology of Ag films as the

temperature changes from room temperature to

700 K were due to the high sensitivity of the

surface diffusion coefficient of Ag adatoms to

temperature. The mobility of Ag adatoms on

Mo(110) was reduced near room temperature and

at a temperature approaching 700 K the diffusion

mean free path was long enough for adatoms to

reach the step edges. Occurrence of mass

transport over step edges in the present work was

not investigated. The desorption process of silver

atoms from Mo(110) at 700 K did not take place,

as shown by Bauer [13]. The authors used AES

to study the annealing behaviour (up to 1050 K)

of thick Ag films on the Mo(110) face. Here, the

AES results showed two temperature regions in

which the Auger signal changed only slightly

with temperature. The authors concluded [13]

that these two small decreases in the silver AES

signal at 700 and 900 K were due to

agglomeration of silver atoms. In addition, at a

sample temperature of T = 1000 K it was

observed that silver desorption from Mo(110)

began[13]. There is also no evidence in the

literature for alloy formation between Ag and a

Mo(110) substrate. Similar behavior for the

Table 1. Structural parameters and surface energies () of Mo, Ni3Al, Pb, Ag, Au, and Sn with respect to Miller

plane orientation reproduced from [44,147-148,154,171-178].


10

Fe/Mo(110) adsorption system has been observed

[181]. These authors studied the growth of iron

on a Mo(110) face by STM. The Fe/Mo(110)

system exhibited the formation of large Fe

islands at room temperature. At temperatures

slightly higher than room temperature the growth

proceeded by incorporation of Fe at Mo step

edges.

2.4. Au/Mo(110)

Figure 5 shows STM images of the Mo(110)

surface with varying Au coverage in order to

illustrate the morphology of the Au layers

deposited on Mo at room temperature. Fig. 5(a)

shows a typical STM image corresponding to a

submonolayer coverage of θ 0.4 ML. The

bright features represent gold islands. These were

one monolayer high islands. The Mo substrate

was still visible between the gold islands. An

analysis of the STM measurements indicated that

for coverage less than 1 ML the distribution of

the gold islands was not random on the Mo

terraces. Namely, from Fig. 5(a) it seemed that

the gold islands grew preferentially along the

]100[ and ]111[ directions. The estimated

average diameter of such gold islands was

approxiamtely 12 5 nm. Fig. 5(b) shows an

STM image corresponding to a monolayer

coverage θ 1.0 ML, as deduced from the

preceeding STM images and deposition time. The

first two-dimensional disordered gold layer was

completed in the manner predicted from the

previous AES, LEED, and TDS experimental

[19, 182-183] and theoretical [184] studies. In

addition, it was not possible to identify further

positions of the gold islands in the first layer

because of the observed per-coalescence effect

[185-192] between θ = 0.55-0.65 ML. In Fig.

5(b), arrows indicate where the second two-

dimensional gold layer started to grow after

completion of the first gold layer. Closer

inspection of Fig. 5(b) when the coverage

increased to 1.2 ML showed that the decoration

of substrate steps by one atomic height Au

islands could be distinguished from other

randomly distributed gold islands on the first

disordered gold layer. For a coverage θ 2 ML,

see Fig. 5(d), a second completed but disordered

gold layer was visible, as expected from literature

Figure 5. STM images of Au deposited on Mo(110) face at T = 300 K: (a) = 0.4 ML (2923 Å × 2923 Å, IT =

0.5 nA, Ubias = 1.0 V). (b) 1.0 ML (2480 Å × 2480 Å, IT = 0.5 nA, Ubias = 1.16 V). (c) 1.2 ML (2494 Å ×

2494 Å, IT = 0.5 nA, Ubias = 1.0 V). (d) 2.0 ML (2500 Å × 2500 Å, IT = 0.5 nA, Ubias = 1.16 V). (e) post

annealed 1.0 ML at T = 800 K (5000 Å × 5000 Å, IT = 0.5 nA, Ubias = 1.4 V). (f) post annealed 1.2 ML

at T = 800 K (2410 Å × 2410 Å, IT = 0.5 nA, Ubias = 1.37 V). The insets in (a), (b), (c), (e) and (f) present

differentiated STM images with enhanced contrast. Adapted from [18]. Copyright (2011) by Elsevier.


11

[182]. The higher surface energy of the Mo

substrate (see Table 1) allowed the Au film to

grow in pseudomorphic registry with the

substrate. This was originally postulated by

Gruza and Gillet [182] using a simulation of a

random adsorption process based on the

following assumptions: (a) every atom impinging

on a free site is adsorbed instantaneously. If the

site is occupied the atom joins either the first

neighbour site or a second layer position, (b)

steric effects can induce displacements of 5%

around an adsorption site, (c) an adsorption site is

inhibited when it is surrounded by at least four

ad-atoms. Based on these assumptions and taking

into account the fact that the distance between

neighbouring sites in the (110) molybdenum

plane (2.73Å) is lower than the atomic diameter

of the gold atom (2.88Å), the distribution of gold

atoms in the saturated first layer show that the

ordered zones are no greater than 20Å. When the

sample was post-annealed to 800 K, the

morphology of the surface changed, as shown in

Figs. 5(e) and (f). At this temperature the

rearrangement of an existing gold film was

observed. These STM results showed similarity

to previous AES studies concerning possibility of

a 2D growth mode [183] between 773 and 973 K.

However, one should remember that post

annealing to a certain temperature (in our case

800 K) and growth at the same temperature, 800

K, cannot be compared directly because of

differing kinetic pathways. This rearrangement

that occrued between the monolayer and two

layer regime gave rise to a non-regular well

separated Au island network, see Fig. 5(e) and

Fig. 5(f) for 1.0 and 1.2 ML, respectively. The

1.0 and 1.2 ML coverages discussed here were

the initial coverages before the post-annealing

procedure. The fact that a complete monolayer

annealed at 800 K lead to island formation is

indicative of gold alloying into the molybdenum

substrate. This agrees well with what is known

from literature [182]; 800 K is below desorption

temperature of gold from the Mo(110) face. In

this publication, the authors used thermal

desorption spectroscopy to study the annealing

behaviour (up to 1550 K) of thick Au films on

Mo(110) surfaces. The TDS results have shown

that the desorption process begins at about 1080

K. Further inspection of the data showed

desorption of a full gold multilayer at T = 1180 K

and of a single monolayer at T = 1300 K. The

analysis of post annealed Figures showed that the

height of the gold islands corresponded to one-

atom high islands, mainly 2.3 Å. From the

presented STM results here it is clear the two-

dimensional gold islands observed after the post-

annealing procedure grew on top of the Au-

Mo(110) surface alloy. The rearrangement of Au

films with change in temperature from room

temperature to 800 K was most likely due to the

high sensitivity of the surface diffusion

coefficient of Au ad-atoms to temperature. The

mobility of Au ad-atoms on Mo(110) was

reduced near room temperature and at a

temperature approaching 800 K the diffusion

mean free path was long enough for ad-atoms to

reach the step edges. The occurrence of mass

transport over step edges in this work was not

observed.

2.5. Sn/Mo(110)

In case of this adsorption system at room

temperature two linear regions of the AES(t) plot

for 186 eV molybdenum can be distinguished,

see Fig. 6(a). The scatter of experimental points

for the tin signal was large because this signal

was low because it was taken with the use of the

retarding field analyser (RFA), and the ratio of

noise to signal was large. It was apparent from

this part of the AES(t) plot that two-dimensional

growth of the first and second layer was

observed. Curve (1) in Fig. 6(a) was calculated

for the Frank – van der Merwe (FM) under the

supposition that at the (hS1; t1) point the first tin

layer was completed. From comparison of our

calculated and experimental AES(t) plots it could

be seen that two-dimensional growth of the

second and third tin layer is also observed.

However, it was impossible to distinguish any

distinct breaking point after deposition time t3 on

the experimentally found AES(t) plot. a shape of

the AES(t) plot suggested Stranski – Krastanov

growth mode after formation of the second layer.

Further we will assume, that the first linear

region of the curve corresponds to the formation

of the first layer (θ = 1 ML) of tin, where 1 ML

of Sn(111) film corresponds to an atomic packing

density of 5.481014 atoms/cm2. Fig. 6(b)-(f)

shows STM images of the Mo(110) surface with

varying Sn coverage in order to illustrate the

morphology of Sn layers deposited on Mo at

room temperature. The insets of Fig. 6(b), (e),

and (f) present corresponding differentiated STM

images with enhanced contrast. Fig. 6(b) shows a

typical STM image corresponding to a

submonolayer coverage of θ 0.7 ML. The

bright features show one monolayer high tin

islands. Such behavior of tin in the initial stage of


12

growth on the Mo(110) surface confirmed

experimental results obtained during our AES

measurements and previous studies [24].

Analysis of the STM measurements indicated

that, for a coverage less than 1 ML, the

distribution of the tin islands was random on the

Mo terraces. This can be seen in Fig. 6(b) where

it appeared that the tin islands do not exhibit

preferential growth directions. Fig. 6(c) shows an

STM image corresponding to a complete

monolayer, i.e. θ 1.0 ML, as deduced from our

AES experiment. The first two-dimensional tin

layer is disordered. For a coverage θ 2 ML, see

Fig. 6(d), a second complete but disordered tin

layer was visible, as expected from literature [24]

and our AES studies. The higher surface energy

of the Mo substrate (and the similar atomic size

of Sn, see Table 1) allowed the Sn film to grow

in pseudomorphic registry with the substrate.

When the sample was post-annealed to 800 K,

the morphology of the surface changed, as shown

in Figs. 6(e) and (f). At this temperature the

rearrangement of an existing tin film was

observed. These STM results were consistent

with previous studies concerning the possibility

of a 2D growth mode [24] between 790 and 1020

K. However, again, one should remember that

post annealing a sample grown at room

temerature to the growth temperature (in our case

800 K) is not a process that can be compared

directly to a growth at that same temperature

because of different kinetic pathways. This

rearrangement that occurred between the

monolayer and bilayer regime gave rise to a non-

regular well separated Sn island network, see Fig.

6(e) and Fig.6(f) for 0.7 and 1.0 ML,

respectively. The 0.7 and 1.0 ML used in this

description are the initial coverages measured

before any post annealing procedure. The average

size of the islands was estimated to be about 10

3 nm. The corrugation changed from 0.4 to 2.4 Å

after post annealing. The fact that a complete

monolayer annealed at 800 K lead to island

formation suggested that tin alloyed into the

molybdenum substrate. It is probable that some

of the tin atoms dissolved into the bulk sample to

form an Sn-Mo alloy. This is expected from

literature [24], where it has been shown that 800

K is below desorption temperature of tin from

Mo(110) face. The authors here studied the

Figure 6. Sn deposition on Mo(110) face at T = 300 K. (a) AES(t) plot of Mo MNN and Sn MNN peak

heights. (1) AES(t) plot calculated for the Frank–van der Merwe growth. STM images at coverages: (b)

0.70 ML (2350 Å × 2350 Å, IT = 0.33 nA, Ubias = 987 mV). (c) 1.0 ML (2300 Å × 2300 Å, IT = 0.2 nA, Ubias

= 1.0 V). (d) 2.00 ML (2500 Å × 2500 Å, IT = 0.73 nA, Ubias = 2.1 V). (e) post annealed 0.7 ML at T =

800 K (1360 Å × 1360 Å, IT = 0.21 nA, Ubias = 790 mV). (f) post annealed 1.0 ML at T = 800 K (1210 Å ×

1210 Å, IT = 0.2 nA, Ubias = 1.0 V). The insets in (b), (e) and (f) present differentiated STM images with

enhanced contrast. Adapted from [20]. Copyright (2011) by Elsevier.


13

annealing behaviour (up to 1550 K) of thick Sn

films on the Mo(110) surface using TDS. It was

shown that the desorption process starts at about

1000 K. Desorption of a tin multilayer was

shown to occur at T = 1180 K and a monolayer at

about T = 1500 K. Analysis of Fig.6 showed, that

the height of the tin islands largely corresponded

to single atomic height islands. It can be deduced

from the STM presented here, that the two-

dimensional tin islands that were observed after

the post annealing procedure most likely grew on

the Sn-Mo(110) surface alloy. It should be noted

that the incorporation of tin into the Mo(ll0)

surface that is proposed here is unusual, because

firstly neither of bulk alloys nor stable

intermetallic compounds are known for this

system, and secondly in other metal systems

where a low surface energy metal is adsorbed on

a high surface energy substrate, the first

monolayer does not mix with the substrate. This

has only been observed previously when a

substrate is soluble in the related film as for Pd

on Mo(ll0) [193], or when the adsorbate and

substrate material form intermetallic compounds

such as Pb on Au(100) [194-195]. On the other

hand, tin is too large for molybdenum to form

stable bulk compounds (high temperature and

pressure is necessary to stabilize such

compound), yet on the surface these size

restrictions are much less dominant. To a certain

extent a somewhat larger size of the adatom can

even be favourable for incorporation because in

this case the surface can achieve a larger packing

density, which is favourable. The rearrangement

of Sn films with change in temperature from

room temperature to 800 K was probably due to

the high sensitivity of the surface diffusion

coefficient of Sn ad-atoms to temperature. The

mobility of Sn ad-atoms on Mo(110) was reduced

near room temperature and increases with

temperature until at temperatures approaching

800 K the diffusion mean free path was long

enough for ad-atoms to reach the step edges. The

occurrence of mass transport over step edges in

this work was not observed. Another important

parameter, which determines the structure of

epitaxial films, is the lattice misfit f [196-199] of

the film to the substrate and can be

accommodated by elastic strain or by the

formation of dislocations. Misfit is given by f =

(aS – aA)/aA, where aS and aA are the lattice

parameters of the substrate and adsorbate,

respectively. Since the film elastic energy is

proportional to f 2, and dislocation energy is

proportional to f, a critical misfit exists below

which pseudomorphic growth can be expected. In

addition, strain energy contributes to the surface

energy of the deposited layer. As a result, if the

strain energy becomes large it could be

energetically more favourable to form three-

dimensional islands on top of the initially formed

film (SK growth mode). Comparison of misfit for

different adsorbates is presented in the Table 1.

2.6. Pb/Mo(110)

Figure 7 shows STM images from the Mo(110)

surface with different Pb coverages in order to

illustrate the morphology of the Pb layers

deposited on Mo at room temperature. A typical

STM image for 0.4 ML Pb deposition can be

seen in Fig. 7(a). Pb islands with an irregular

shape are shown as bright features and were

deduced to be monolayer height islands. As the

Pb coverage approached 1 ML, the Pb wet the

Mo(110) surface almost completely, as can be

seen in Fig.7(b). It is not easy to determine

whether an adsorbate wets a surface or not by

STM. However, almost perfect wetting occurred

due to the high specific surface free energy of the

Mo(110) surface compared to that of the Pb(111)

surface. As the total specific surface free energy

is always minimized by nature, a covered

Mo(110) surface was favoured. As it can be seen

from the inset of Fig.7(b), the first Pb layer on

the Mo(110) surface was not free from

dislocation defects. The pseudomorphic growth

of the first two-dimensional Pb layer was

thermodynamically driven by the lower surface

free energy of Pb, and is the dominant process so

the Pb layer wets the Mo surface despite the

elastic energy required as a result of the large

lattice mismatch. Detailed inspection of Fig.7(b)

showed that the second lead layer starts to grow

before complete formation of the first layer. The

arrows in Fig.7(b) indicate where the Pb adatoms

of the second layer started to nucleate [200]

which occurred preferentially at dislocation

defects of the first Pb layer. As it can be seen, the

density of the dislocation defects was higher at

monatomic ledges. At higher coverages (1 ML <

5 ML) the growth became three-dimensional,

(Fig.7(c)-(f)). These results contradict the

literature data [25], where layer-by-layer growth

at room temperature was reported. Fig.7(c) and

(e) show STM images after 1.2 and 2.5 ML of Pb

were deposited on the Mo(110) surface at RT.

Height profiles that were taken across the Pb

islands (Fig.7(d)) show that the Pb crystallites

had a thickness of about 6 Å. The interlayer


14

spacing along the [111] direction of Pb was d111 =

2.86 Å and therefore the islands had a thickness

corresponding to the stacking of two atomic

layers. The heights of Pb islands were measured

from the wetting layer. At coverages below 2

ML, Pb islands were observed with an average

diameter of 800 Å and a standard deviation of 70

Å. As the Pb coverage was increased up to 2.5

ML (Fig.7(e)), more than 75% of the islands

were two layers in height. Four-layer high islands

were also observed, however, but they

represented less than 25% of the total population

of the Pb islands. In Fig.7(e), the heights of the

Pb islands measured in atomic layers are

indicated. The observed Pb islands on the wetting

layer had regular shapes, sharp edges and smooth

tops, and tended to grow up with higher

coverages suggesting that such islands are stable.

It should be pointed out that above > 1 ML,

one-layer thick Pb islands were never observed.

At coverages greater than 3 ML (Fig.7(f)), further

formation of islands with sharp edges and smooth

tops was observed. The island height distribution

(number of Pb islands with a given height),

however, showed strong peaks at relative heights

corresponding to the number (N = 2, 4, and 6) of

Pb atomic layers. During post-annealing to 1050

K, a change of the height of the lead ‘magic

islands’ was observed. Fig.8 shows STM images

obtained after the post-deposition annealing

(1050 K, 5 sec.) of Pb films ( = 5.0 ML)

deposited on Mo(110) at room temperature [20,

201]. Height profile (C), taken across the Pb

islands (Fig. 8(c)), show that only islands

containing 2, 5, and 7 atomic layers of Pb were

observed above the wetting layer. Close

inspection of the STM topography image in

Fig.7(b) revealed the presence of two well-

ordered lead superstructures denoted 1 and 2

[23]. Their geometrical parameters are described

Figure 7. STM images of Pb deposited on Mo(110) face at T = 300 K: (a) = 0.4 ML (1140 Å × 1140 Å, IT =

0.2 nA, Ubias = 1 V). (b) = 1.0 ML (1951 Å × 1951 Å, IT = 0.3 nA, Ubias = 800 mV). The inset presents zoom

(242 Å × 242 Å) in the area indicated by a square. Arrows indicate the places where the Pb atoms of the

second layer start to nucleate at dislocations of the first Pb layer. (c) = 1.2 ML (10000 Å × 10000 Å, IT = 0.2

nA, Ubias = 1 V). (d) Height profile along the line B from the image in (c) demonstrating that the height of the

islands corresponds to the height of a double Pb step height. (e) = 2.5 ML (8563 Å × 8563 Å, IT = 0.2 nA,

Ubias = 1.0 V). The height of the Pb islands corresponds to 2 and 4 atomic layers of Pb(111). (f) = 5 ML

(10000 Å × 10000 Å, IT = 0.2 nA, Ubias = 1 V). Island height distribution (number of Pb islands with a given

height) obtained from this image has showed the strongly preferred heights of the Pb islands corresponding to

2, 4 and 6 atomic layers of Pb(111) [22]. Copyright (2009) by The American Physical Society.


15

in the caption of Fig.9(a)-(b). The structures 1

and 2 were observed simultaneously in different

areas of the Mo(110) substrate. STM images

indicated in both cases that a long-range order in

the surface system was present. The density of Pb

atoms determined for the 1 and 2 structures

were 4.5×1014 and 6.4×1014 atoms/cm2,

respectively, which corresponds to 48% and 68%

of the density of the close-packed (111) Pb layer

(equal to 9.4×1014 atoms/cm2). Taking into

Figure 8. STM images of Pb deposited (Pb 5.0 ML) on Mo(110) face at T = 300 K and post annealed at 1050

K: (a) (20000 Å × 20000 Å, IT = 0.21 nA, Ubias = 559 mV) showing the observed ring morphology (indicated by

white arrows) on the 5 atomic layers of Pb(111) islands. (b) (6100 Å × 6100 Å, IT = 1.18 nA, Ubias = 298 mV).

Height of the islands corresponds to 2, 5 and 7 atomic layers of Pb(111) [201]. (c) Height profile along line C

from the image in (b) demonstrating that the heights of the rings corresponding to 2 atomic layers of Pb(111). (a)

is adapted from [20]. Copyright (2011) by Elsevier.

Figure 9. Pb deposition on Mo(110) face at T = 300 K. STM images (differentiated in insets) of two kinds of

Pb superstructures: (a) 1 (80 Å × 80 Å, IT = 1.7 nA, Ubias = 149 mV). (b) 2 (50 Å × 50 Å, IT = 3.3 nA, Ubias =

76 mV). White parallelograms indicate unit cells of the respective Pb structures with geometrical parameters

measured to be equal to a = 4.3 Å, b =5.2 Å, and α = 84° for 1, and a = 4.1 Å, b = 3.8 Å, and α = 92° for 2.

(c) and (d) presents overlap of the geometrical models of the superstructure 1 and 2 with its filtered STM

images, respectively. Parameters of the simulated unit cells are a = 4.1 Å, b = 5.2 Å, α = 84° and a = 4.1 Å, b

= 3.9 Å, α = 90° for 1 and 2, respectively. Red dashed lines denote the supercells used in simulations [23].

Copyright (2009) by The American Physical Society.


16

account the experimental results, structural

models of the Pb/Mo(110) surface have been

constructed reproducing the topography of the

obtained STM images. The lateral geometrical

properties of this model (distances between the

nearest lead adatoms along the directions a⃗ and b⃗ and the angle between them) are almost the same

as those deduced from the STM measurements

(cf. captions of Figs. 9(a) and (c)). Structural

relaxation showed that such a model would be

stable, i.e., the lateral positions of all lead

adatoms in the relaxed structure were the same as

in the starting configuration while PbS adatoms

protruded 0.08 Å above PbL. The superstructure

2 is, in turn, modeled by an adlayer with PbL,

PbS, and PbT atoms (cf. Fig.9(d), corresponding

to the coverage of 4/9 ML. The lateral parameters

of this model (i.e., a⃗ , b⃗ and α) reproduce well the

STM data (cf. captions of Figs.9(b) and (d)).

Results of structural simulations show PbT

adatoms located 0.09 Å above PbS and 0.17 Å

above PbL while the lateral positions of all lead

atoms remain fixed during the relaxation. The

supercells used for simulation of both

superstructures are different, and therefore, a

direct comparison of their total energies is not

possible. Thus, to determine the energy

relationship between them, the average

adsorption energy (Eb) per one Pb adatom has

been calulated defined as Eb =

−1

N(EPb/Mo(110) − EMo(110) − NEPb ), where

EMo(110) is the energy of a clean Mo(110) slab,

EPb/Mo(110) is the energy of an adsorbed slab, and

EPb is the energy of free Pb atom, and N is the

number of adsorbate atoms in the cell. To this

end, Eb was averaged over two or four adatoms

from the supercell corresponding to the

superstructures 1 and 2, respectively (cf.

Figs.9(c) and (d)). Obtained averaged adsorption

energies are equal to 4.28 eV for 1 and 4.29 for

2. It should be noted that the energetic properties

of the lead superstructures considered in this

section are consistent with experimental results

presented in Ref. [24], indicating that adsorption

energy of lead adatoms increases with coverage.

Indeed, adsorption energies calculated for dense

superstructures corresponding to a coverage of

2/3 ML were above 4.4 eV, and are noticeably

higher than the values (below 4.3 eV) obtained

for the more dilute superstructures given in

Figs.9(c) and (d).

2.7. Pb/Ni3Al(111)

In the growth of lead on Ni3Al(111) at room

temperature, one linear component of the AES(t)

plots for both 61 eV nickel and 94 eV lead peaks

were observed, see Fig. 10(a). This component of

the AES(t) plots indicated a two-dimensional

(2D) growth of the first lead layer. After the first

monolayer, the quasi-linear shape of the second

region of the AES(t) suggested that a two-

dimensional growth continued that was not of the

standard Frank – van der Merwe (FM) type. In

this region the calculated AES(t) plot for a

substrate with FM growth did not fit the

experimental data (see curve (1) in Fig. 10(a)). A

much better fit was obtained for the plots for

growth up to 2.5 ML when a growth of flat Pb

islands composed of three atomic layers (curve

(2) in Fig. 10(a)) were assumed in the

calculations. Another assumption that was made

was that the first linear region of the curve

corresponded to the formation of the first layer (θ

= 1.0 ML) of lead. Fig. 10(b)-(d) show STM

images of Ni3Al(111) surfaces with various Pb

coverages and illustrate the evolution of growth

morphology of the Pb layer deposited on

Ni3Al(111) at room temperature. Fig. 10(b)

shows a typical STM image corresponding to a

monolayer coverage θ = 1.0 ML where Pb was

seen to wet the Ni3Al(111) surface almost

completely. At this coverage only LEED spots

associated with the p(4×4) structure were

observed [32]. The density of Pb atoms in the

p(4×4) structure is known to be 10.42×1014

atom/cm2, which corresponds to 110% of the

density of the close-packed (111) Pb layer. The

almost perfect wetting in this system was due to

the high specific surface free energy of the

Ni3Al(111) surface as compared with that of the

Pb(111) surface. Since the total specific surface

free energy is always required to be minimised, a

covered Ni3Al(111) surface was favoured. As

seen in the inset of Fig. 10(b), the first Pb layer

on the Ni3Al(111) surface was not free from

dislocation defects. The arrows in Fig. 10(b)

indicate areas where holes with an average depth

of dh = 1.02 ± 0.35 Å in the wetting layer were

observed. The observed two-dimensional growth

of the first Pb wetting layer was

thermodynamically driven by the lower surface

free energy of Pb compared to Ni3Al, see Table

1. This wetting was still favorable despite the

elastic energy resulting from the large lattice

mismatch between Pb and Ni3Al, ~39% [32]. A

similar wetting layer growth with dislocations

was observed for Pb adsorption on Mo(110) [22]

as described above in Section 2.5. The presence


17

of compressed wetting layers with a density

higher than the metallic Pb(111) density by

approximately 5% has been identified as the key

factor that provides unusually fast and non-

classical kinetics in the uniform height islands

[34,53]. It is not unusual to expect that the p(44)

phase plays the same role and superfast diffusion

could have been present in this system as well. At

coverages of > 1 ML, three-dimensional growth

was observed. Figure 10(c) shows STM images

taken after 3.5 ML of Pb was deposited on the

Ni3Al(111) surface at room temperature. The Pb

islands have a hexagonal shape, flat top and

thickness of approximately 9, 14, 20, or 32 Å.

The interlayer spacing along the [111] direction

of Pb was measured to be d111 = 2.86 Å and

therefore the islands had a thickness

corresponding to 3, 5, 7 or 11 atomic layers,

respectively. As Pb coverage was increased to 5.5

ML, three types of regular hexagonal islands with

length side 25, 30 and 45 nm were measured

(Fig. 10(d)). Smaller irregularly shaped Pb

islands were also observed between the main

regular hexagonal islands. Furthermore, three

different adsorption configurations of Pb were

distinguishable, rotated by α = 10° ± 1°and β =

30° ± 6° with respect to each other. It should be

noted that the different orientation of these

hexagonal Pb islands was not correlated with

their size. The presence of these multilayer lead

islands agreed with the island growth deduced

from AES analysis at room temperature [32].

During post-deposition annealing (T = 400 K, 60

sec.) of Pb films ( = 5.5 ML), a change of the

height of the lead ‘magic islands’ was observed

as seen in Fig. 10(e)-(f). Only islands containing

Figure 10. Pb deposition on Ni3Al(111) face at T = 300 K. (a) AES(t) plot of Ni MVV and Pb NVV peak

heights. (1) – AES(t) plot calculated for Frank–van der Merwe growth. (2) – AES(t) plot calculated for the

growth of three atomic layer thickness flat islands on the first lead layer. STM images at coverages: (b) =

1.00 ML (2070 Å × 2070 Å, IT = 122 pA, Ubias = 2.0 V); The inset presents zoom (392 Å × 392 Å) in the area

marked by a square. Arrows indicate holes in the wetting layer. (c) = 3.5 ML (1520 Å × 1520 Å, IT = 100

pA, Ubias = 3.2 V). The height of the Pb islands corresponds to 3, 5, 7 and 11 atomic layers of Pb(111).

Hexagonal shape of Pb islands is marked. (d) = 5.5 ML (4400 Å × 4400 Å, IT = 146 pA, Ubias = 7.1 V). The

Pb islands have regular hexagonal shape I, II and III with the length of the side 25, 30 and 45 nm,

respectively. The different orientation of the Pb islands with respect to each other is marked. (e) post annealed

= 5.5 ML at T = 400 K (6600 Å × 6600 Å, IT = 130 pA, Ubias = 7.6 V). The height of the Pb islands

corresponds to 2 atomic layers of Pb(111) measured from the wetting layer. The inset present differentiated

STM images with enhanced contrast. (f) Height profile along line D from the image in (e). Adapted from [33].

Copyright (2013) by Elsevier.


18

two atomic layers of Pb are observed above the

wetting layer. LEED patterns only corresponding

to a weakly ordered p(44)-Pb structure were

observed after post-deposition annealing [32].

These STM results confirmed AES studies at T =

400 K concerning a 3DIM growth mode at the

same height above the wetting layer. Table 2

shows the addition of the Pb-Mo(110) and Pb-

Ni3Al(111) adsorption systems to the known

group of systems where uniform island height

selection during metal thin film growth has

already been observed and interpreted in terms of

quantum size effects (QSE) [60-65]. Typically,

these electronic effects have been observed on

semiconducting [41,43-59] or metal substrates

[22, 32-42,58]. For example, STM/STS

observations of Pb islands on the Cu(111) surface

face [40] indicated that the equilibrium

distribution of island heights showed some

heights appearing much more frequently than

others. In all these systems a confining barrier

restricted the electron to be within the film

because of misaligned electronic bands in the

metal film and the substrate so no states were

available to move from the film to the substrate.

‘Magic’ preferred heights correspond to islands

with a quantum well state far from the Fermi

energy, while the ‘forbidden’ heights appear to

be those that have a quantum well state close to

the Fermi level. Another typical example of

uniform island growth is Pb deposited on Si(111)

at a temperature between 170 and 250 K and

typically results in the formation of Pb islands

with N = 4(unstable), 5, 6, 7 and 9 atomic layer

heights above the wetting layer. Nevertheless, the

occurrence of ‘magic heights’ is related to the

increased stability associated with islands of

specific thickness. Ayuela and co-workers [202]

studied quantum size effects of Pb overlayers

using density-functional theory (DFT)

calculations and analytical models. This work

demonstrated that the high stability of Pb islands

on metallic or semiconductor substrates at even

higher coverages was supported by an extra

second quantum beat pattern in the energetics of

the metal film as a function of the number of

Table 2. The table summarizes the observed layer thicknesses (N) of ‘magic height’ Pb islands for different

adsorption systems. Layer thicknesses are measured with respect to ‘wetting layer.’ (u) – describes unstable Pb

islands. Reproduced from [20,22,32-36,40,43-48,50,52,61,201].


19

atomic layers. It seemed that this pattern was

triggered by the butterfly-like shape of the Fermi

surface of lead in the [111] direction and

supported the detection of stable ‘magic islands’

of greater heights than measured until now. The

most stable ‘magic islands’ from this DFT study

had heights corresponding to the number (N = 2,

4, 6, 7, 9, 11, 13, 14, 16, 18, 20 …) of lead

atomic layers. Recently, Özer et al. [203] and Jia

at al [204] have shown that, similar to the case of

pure Pb, the Pb1-xBix alloy films on Si(111) 7×7

exhibit a reentrant bilayer-by-bilayer growth

mode but with different quantum beating

patterns. The observed stable thickness sequence

was 4, 6*7, 9, 11, 13, 15, 17, 19*20…ML, where

the asterisks indicate the locations of the even-

odd crossover. At the higher concentration of

Pb80Bi20, the alloys did not exhibit quantum

growth but simply followed classical layer-by-

layer growth.

3. Molecules and metal growth on the

Al2O3/Ni3Al(111)

3.1. Al2O3/Ni3Al(111)

It is known that the controlled oxidation of the

clean Ni3Al(111) surface results in the formation

of a closed double layer alumina film with a

distinctive long-range order [156-157, 205-206].

Recent SPA-LEED and STM investigations [92]

showed that under the preparation conditions

used, the Al2O3 film forms a long-range ordered

hexagonal superstructure on the nanometer scale.

Fig. 11(a) shows an overview SPA-LEED image

of the alumina film on Ni3Al(111) at an electron

energy of 110 eV. Apart from the aluminum

p(22) superstructure of the substrate

(Fig. 11(b)), a very large number of additional

superstructure spots originating from the Al2O3

film could be seen. Around the (0,0) spot two

rings consisting of twelve superstructure spots

each (Fig. 11(a)) could be observed. These two

rings could be related to two superstructures, the

“network” and the “dot” structure. The fact, that

we observed 12 spots, was related to the

superposition of two hexagonal rotational

domains of the alumina film, which have already

been observed by STM [207]. The angle between

domains was 24° which was in good agreement

with the STM results. The larger of the two rings

could be attributed to the diffraction from the

“network” structure. The size of the unit cell of

this structure (blue lines in Fig. 11(b)) could be

calculated when using the known size of the

substrate lattice. The value of bnet(SPA-LEED) is

24 Å. Similarly, the inner circle could be

attributed to diffraction from the “dot” structure.

In this case the value for the unit cell size

bdot(SPA-LEED) was equal to 41.6 Å.

Furthermore, it was determined that the angle of

rotation between the network structure and the

substrate was 18° as calculated from the SPA-

LEED image. Based on these findings it was

possible to simulate the LEED pattern using the

substrate lattice and unit cells of the “network”

and “dot” structures, blue and red circles in

Figs.11(b), respectively. The strong dependence

of the corrugation on the tunneling voltage can be

seen in the STM images of Fig. 12. At bias

Figure 11. (a) SPA-LEED image of the Al2O3/Ni3Al(111) taken at an electron energy of 110 eV. (b) Simulation

of the LEED image. The units cells are: p(2×2) aluminum sub-lattice of the substrate (green), network structure

(blue), and dot structure (red). Adapted from [92]. Copyright (2005) by Elsevier.


20

voltages around Ubias = 3.2 V (see Fig. 12(a)) a

hexagonal arrangement of hollows was observed

where every hollow was surrounded by a

hexagonal ring of bright dots. This structure

possessed an apparent lattice constant of 24 Å

and will be referred to as ‘‘network structure’’ in

the following text. However, closer inspection of

this structure indicated that not all hollows have

the same depth. Some hollows were up to 40 pm

shallower than the others. This fact is clearly

visible in the cross section shown in Fig.12(d).

The shallower hollows formed a regular

arrangement, which formed the basis of a

hexagonal unit cell with a lattice vector of 41.6

Å. This unit cell of the alumina film was more

easily observed in the STM images, Fig. 12(b)

and (c), due to a contrast reversal, which was

found at a bias voltage of 2.0 and 4.2 V, and will

be referred to as ‘‘dot structure’’ in the following

text. Recent AFM measurements have shown

that, indeed, the surface was atomically flat and

that the ‘‘dot structure’’ corresponded to the

actual surface unit cell [208-209]. Scanning

tunneling spectroscopy (STS) was used to study

the electronic properties of the system

Al2O3/Ni3Al(111) [157]. In Fig. 13(a) STS

spectra are shown together with the

corresponding STM image in which the positions

are shown where the spectra were acquired.

Curves (A) and (B) were taken at a deep hollow

site and on the hexagon surrounding that site,

respectively. In the range -4 V to -6 V the main

ascent of the density of states of the oxide

valence band was visible. At the large tip-sample

separation used for these spectra, no electronic

states in the range -4 V to 2.7 V were detected.

Above a value of 2.7 V an increase of the density

of states was seen, which corresponded to the

lower edge of the oxide conduction band. From

this it was possible to derive a value of 6.7 eV for

the band gap of the Al2O3 film. The band gap of

an Al2O3 film prepared on NiAl(110) has also

been quoted before at 6.7eV [210]. The

comparison of curves (A) and (B) in Fig. 13(a)

shows that the band gap measured by STS was a

function of the tip position. In Fig. 13(b) an STM

image is shown for which a set of tunneling

spectra were measured at different locations (A)-

(F) and the corresponding STS spectra are

displayed in Fig. 13(c). The centre of this STM

Figure 12. LT-STM images of the Al2O3/Ni3Al(111) measured at T = 23 K and different bias voltages. The

unit cell of the hexagonal alumina superstructure is drawn in red. The image parameters are: (a) The ‘‘network

structure’’ at Ubias = 3.2 V (278 Å × 278 Å, IT = 122 pA), (b) The ‘‘dot structure’’ at Ubias = 2.0 V (278 Å × 278

Å, IT = 105 pA), (c) The ‘‘dot structure’’ with inverted contrast at Ubias = 4.2 V (123 Å × 123 Å, IT = 103 pA).

In (d) two cross sections taken at the positions indicated in (a) and (b) are shown. (a), (b) and (d) are adapted

from [92]. Copyright (2005) by Elsevier. (c) is adapted from [95]. Copyright (2008) by Elsevier.


21

image (A) shows a shallow hollow that was a

coincidence point of the commensurate Al2O3

structure [92]. The spectra were taken over a bias

voltage range 1.5 V to 3.3 V. This corresponded

to an energy range in the vicinity of the oxide

conduction band edge where electronic states

originating from the oxide could be detected.

Spectrum (A) was measured directly at the

coincidence point. It exhibited one maximum at

approximately 2.4 V, which was in agreement

with previous studies [211]. Spectra (A) to (F)

were measured at an increasingly larger distance

from the coincidence point. Going from spectrum

(B) to spectrum (E) the maximum at 2.4 V

decreased and was no longer visible in spectrum

(E). In addition an electronic state around 3.0 V

appears. In Spectra (B) and (C) the maximum

around 2.4 V could be distinguished as well as

the one around 3.0 V. These results clearly

showed that in this energy range there existed

two electronic states with different origins at

different locations within the oxide supercell. The

electronic state at 3.0 V was present everywhere

on the surface except the coincidence points.

Spectrum (F) was measured in the centre of a

deeper hollow of the structure and was nearly

identical to spectrum (E), which corresponded to

a completely different point on the alumina film.

Previous studies on the Al2O3/Ni3Al(111) system

showed a pronounced template effect for growth

of a number of metals. While the ‘‘network

structure’’ was shown to have a direct influence

on the growth of V [212] and Mn [213] clusters,

a template effect of the ‘‘dot structure’’ for Cu

[212], Ag [214], Au [214] and Pd [107] clusters

has been demonstrated. The next logical step was

the investigation of less strongly interacting

adsorbates such as organic molecules and will be

discussed in the next section.

3.2. C32H16N8Cu/Al2O3/Ni3Al(111)

Thin films formed from organic molecules

including Copper phthalocyanine are rapidly

gaining traction in different technological fields

[215]. Copper phthalocyanine is an organic

molecule with a molecular weight of 576 g/mol

and a side length of 12 Å. It possesses a D4h-

symmetry and is completely planar in the gas

phase [216]. The extensive aromatic -electron

system and the chelating properties of the

phthalocyanine core result in stable complexes

with copper and many other metal ions. In the

ground state copper is in the oxidation state +II.

Pure CuPc is known to crystallize in several

modifications [217]. The - and -modification

are the most stable structures. Both are composed

of stacked layers of planar Pc molecules, which

show different stacking sequences and layer

spacings of 0.34 nm [216] and 0.33 nm [218],

respectively. The stacked arrangement of the

molecules enables strong interactions between

their p-electron systems (‘‘-stacking’’). This

results in a high thermal stability and low vapor

pressure. DFT calculations have shown that the

HOMO is strongly affected by the dx2

-y2 -orbital

of the copper ion [219]. This results in a

relatively high energy due to the electronic

splitting in the quadratic planar ligand field.

Because of the d9-configuration of the copper ion

the HOMO is only half filled and may

theoretically take part in the tunneling process.

However, due to the missing z-component of the

dx2

-y2 –orbital this seems to be improbable so the

Figure 13. LT-STS measurements of the Al2O3/Ni3Al(111) at T = 23 K [157]. (a) Tunneling spectra measured at

the points indicated on the STM image (50 Å × 50 Å, IT = 100 pA) in the inset. Curves (A) and (B) were

acquired after the feedback loop was opened at 3.2 V. The corresponding modulation frequency was 10 kHz

with an amplitude of 40 mV. (b) Fourier filtered STM image (53 Å × 53 Å, IT = 95 pA, Ubias = 3.2 V). (c)

Tunneling spectra (A) - (F) acquired at the points indicated in (b).


22

STM depiction should be dominated by the

LUMO, which has eg-symmetry. The big energy

gap of more than 1.5 eV between the LUMO and

the higher unoccupied orbitals should prevent

their participation in the tunneling process. This

assumption was proven by our STM

investigations [95] at a coverage of CuPc = 0.03

0.02 ML, Fig. 14. Individual molecules were

imaged as four-lobe protrusions, which was in

good accordance with the calculated LUMO

[220]. Small differences like the obviously

missing electron density at the outer carbon

atoms of the benzene rings were most likely due

to interactions with the substrate. The

surprisingly large side length of the adsorbed

CuPc in the STM images of ~ 20 Å compared to

the expected value of 12 Å (see Fig. 14(a)) was a

frequently observed phenomenon, which could

be explained by a small overlap of the orbitals of

tip and molecule even at greater distances [221].

This resulted in a small tunneling current, even

when the tip was not exactly positioned above the

molecule. At this coverage of CuPc = 0.03 0.02

ML local dI/dV curves of individual CuPc

molecules were recorded. Since the CuPc

molecules and their arrangement on the surface

were quite sensitive to the alternating high

electric field of the tip during the STS

measurements some precautions had to be taken.

Only spectra that did not influence the position

and the shape of the molecules were considered.

To ensure this, STM images of the individual

molecule were taken before and after the

spectroscopic measurements. Interestingly at

different positions of the tip above the molecule

(Fig.14(c)) the local dI/dV curves shown in

Fig.14(d) looked quite similar. They exhibit only

one maximum centered around 1.2 eV, which

could be clearly identified as the energy of the

degenerate LUMO. This exhibited a marked

difference to former investigations of CuPc

adsorbed on Al2O3/NiAl(110) by Qiu et al. [113].

These authors observed two distinct molecular

symmetries (D4h and D2h). The ones with D2h-

symmetry featured two different maxima in the

local dI/dV curves, which were detected on two

pairs of oppositely located lobes. This was

explained by the lifting of the degeneracy of the

LUMO by reason of a Jahn–Teller effect.

However, we could only detect molecules with

D4h-symmetry and degenerate LUMO on

Al2O3/Ni3Al(111). This could have been due to

small differences in the interaction of CuPc with

Al2O3/N3Al(111) as compared to

Figure 14. Copper phthalocyanine (C32H16N8Cu). (a) Chemical structure. (b) Calculated shape of the LUMO

[109]. (c) STM image measured at T = 23 K of a single CuPc molecule adsorbed on Al2O3/Ni3Al(111) (33 Å ×

33 Å, IT = 100 pA, Ubias = 1.2 V). (d) dI/dV tunneling spectra corresponding to the numbered positions in (c)

(AC modulation voltage of 60 mV at 1 kHz). Adapted from [95]. Copyright (2008) by Elsevier.


23

Al2O3/NiAl(110). In fact, we found only one

molecular CuPc species, which adopted two

molecular orientations rotated by 30 with

respect to each other. Fig.15(a) shows larger

STM image of the oxidized Ni3Al(111) surface

with a CuPc coverage of CuPc = 0.03 0.02 ML.

At this coverage the molecules were always

adsorbed with the molecular plane parallel to the

substrate surface. At this low coverage only

individual molecules were found, which

apparently did not form an ordered long-range

structure. Closer inspection of the images showed

that the metal centre of the CuPc was always

positioned above one of the bright spots of the

hexagons surrounding the hollows of the

‘‘network structure.’’ However, the ‘‘dot

structure’’ also seemed to effect the position of

the molecules. Even though they were never

centered directly on a coincidence point [92] in

this structure. A significant affinity to surface

defects, e.g. terrace edges or imperfections in the

oxide layer, was not, however, observed at low

temperature and this was probably due to the low

mobility of the CuPc. Furthermore, a distinct

orientation of the CuPc molecules with respect to

the substrate could be observed at low coverage.

As demonstrated in Fig. 15(a) only two different

molecular adsorption configurations of CuPc

were distinguishable, rotated by 30 with respect

to each other. This suggested that a distinct

adsorption geometry was preferred in order to

maximize the adsorption energy as a result of the

interaction of the square shaped molecule and the

hexagonal alumina film. The intermolecular

interactions between the aromatic -systems of

the molecules became ever more important with

increasing CuPc coverage. Fig.15(b) shows the

template effect of the substrate was not strong

enough to compensate for these interactions, even

at a moderate coverage of CuPc = 1.0 0.2 ML.

At 140 K the growth of the second molecular

layer started before the monolayer was complete,

resulting in the formation of randomly orientated

molecular clusters [95]. Here, the molecules were

always oriented with the molecular plane parallel

to the substrate. Interestingly, even in the second

layer the CuPc could be imaged with

submolecular resolution. In an earlier work by

Lippel et al. [109] it was reported that this was

not possible on Cu(100), Si(111) and Au(111).

The measured height difference between

molecules in the first and second layer of 0.35

nm (Fig.15(c)) was comparable to the layer

distance of 0.34 nm in the -modification of pure

Figure 15. LT-STM images measured at T = 23 K of CuPc adsorbed on Al2O3/Ni3Al(111). (a) = 0.03 0.02

ML (130 Å × 130 Å, IT = 107 pA, Ubias = 1.6 V). The different orientation of the single molecules with respect

to each other is marked. (b) = 1.0 0.2 ML (174 Å × 174 Å, IT = 110 pA, Ubias = 3.2 V). (c) Height profile

along the marked line in (b). (d) = 0.03 0.02 ML (138 Å × 138 Å, IT = 110 pA, Ubias = 3.2 V), CuPc on the

“striped phase” of Al2O3. [222]. (a)-(c) are taken from [95]. Copyright (2008) by Elsevier.


24

CuPc, indicating a similar intermolecular

interaction as in the bulk of this compound. The

layer distance, under normal thermodynamic

conditions, in the stable -modification was 0.33

nm. Figures 15(d) shows a STM image of the

CuPc molecules adsorbed on the “striped phase”

of Al2O3 film [222]. In order to characterize

interaction of the molecules and the surface in

greater detail the surface was annealed. After

annealing to 300 K the formation of CuPc

agglomerates could be observed, indicating a

high lateral mobility of CuPc. Single molecules

could not be observed. After annealing the

sample to 350 K the CuPc was desorbed

completely. According to this, it could be

concluded that CuPc only weakly chemisorbs on

to the oxidized Ni3Al(111) surface and that the

intermolecular forces are much stronger than the

molecule–surface bond [95].

3.3. Fe clusters on Al2O3/Ni3Al(111)

Several recent review articles have described the

growth, structure and electronic properties of

various metallic overlayers on different oxide

surfaces [115, 223-228]. In this section, we

concentrate our review on recent results

concerning the behavior of Fe on Al2O3. The iron

clusters grew and gradually covered the entire

Al2O3/Ni3Al(111) surface with increasing

coverage [96]. However, the growth of a

continuous wetting layer of Fe has not been

observed. This result was in agreement with the

thermodynamic criterion for the wetting of

alumina [229-231]. From energetic

considerations wetting should not occur on the

alumina surface if Eadh < 2 v/m, where Eadh is the

adhesion energy (?) and v/m is the surface free

energy at the vacuum metal interface for liquid

metal. The definition of Eadh is the energy needed

to separate the metal and the oxide at their

interface resulting in a bare oxide surface. Eadh

for iron on alumina was measured to be 1205

mJ/m2 at 1853 K [231]. Eadh is mainly energetic

and therefore nearly temperature independent.

v/m could be extrapolated to temperatures below

the phase transition entailing an increasing v/m

by measuring v/m at 1853 K to be 1787 mJ/m2

[229-231]. The thermodynamic criterion (no

wetting) was consistent with the observation of

Figure 16. STM images of Fe clusters on Al2O3/Ni3Al(111): (a) deposited at 160 K ( = 0.07 ± 0.02 ML,

198 Å × 98 Å, IT = 106 pA, Ubias = 3.5 and 4.2 V in upper and lower part, respectively; TSTM = 23 K). The unit

cell of the hexagonal alumina superstructure is drawn in red. (b) Corresponding line profiles along the blue and

green lines in (a). (c) deposited at 130 K ( = 1.0 ± 0.2 ML, 681 Å × 681 Å, IT = 78 pA, Ubias = 0.7 V,

TSTM = 300 K). The FFT shown in the inset evidences a hexagonal arrangement of the iron clusters with a

nearest neighbor distance of 24 Å. (d) Fourier filtered STM image (62 Å × 62 Å, IT = 151 pA, Ubias = 3.2 V).

(e) Fourier filtered STM image (62 Å × 62 Å, IT = 151 pA, Ubias = 0.19 V), The (22) unit cell of the substrate

is drawn in yellow [222]. (a)-(c) are adapted from [96]. Copyright (2006) by Elsevier.


25

three-dimensional (3D) particles at 300 K. In Fig.

16(a) and (c) the iron clusters grown at a low

deposition temperature (130 K – 160 K) are

depicted. The surface structure grown to low

iron coverage ( = 0.07 ± 0.02 ML) is shown in

Fig. 16(a), (d) and (e). The iron clusters appear as

bright spots on top of the long-range modulated

alumina film. As evidenced in the upper part in

Fig.16(a) taken at Ubias = 3.5 V and in Fig.16(d)

taken at 3.2 V, the iron clusters nucleated first in

the hollows corresponding to points of the

“network structure” of the alumina film. A

similar behavior was observed for V and Mn

[106], also on alumina. The noble metals Ag, Cu,

Au and Pd, however, nucleate preferentially on

the “dot structure” of Al2O3/Ni3Al(111) with

clusters being 41.6 Å apart from each other [106-

107]. The apparent height of the iron clusters

grown between 130 K and 160 K was 2.6 ± 0.3 Å

with respect to the oxide surface, as inferred from

the line profile shown in Fig.16(b). It is

concluded from STM that the islands are one

atomic layer in height. Cluster growth is a non-

equilibrium process and the final macroscopic

state of a system depends on the different

activation energies for all microscopic kinetic

processes that are involved. At 130 K, the growth

was limited to 2D, since a diffusing Fe atom

laterally approaching a 2D island would attach

itself to the step but cannot climb up to reach the

islands top surface. However, a Fe atom landing

on-top of an Fe island can overcome the Ehrlich-

Schwoebel barrier and would therefore move

down-step. At a Fe coverage of = 1.0 ± 0.2 ML

the nucleation sites of the “network structure”

were all filled with Fe clusters which partly

coalesce (Fig.16(c)). From the Fourier transform

(FFT) shown in the inset, a long-range hexagonal

order of the iron clusters that corresponded to a

distance of 24 Å between the clusters could be

observed. Here, the iron clusters were still two-

dimensional and could be deduced from their

Figure 17. (a) Set of STM images of Fe clusters deposited at 160 K on the Al2O3/Ni3Al(111) surface recorded

as a function of bias voltage ( = 0.07 ± 0.02 ML, 480 Å × 480 Å, IT = 102 pA, TSTM = 23 K). (b) and (c)

presents bias voltage dependence of the height and diameter of the Fe clusters marked in the STM images,

respectively. The increase in cluster height at ~ 1.5 eV can be attributed to empty states of the Fe clusters

whereas empty states of the Al2O3 film contribute around ~ 2.4 eV [222].


26

apparent height of 2.6 ± 0.3 Å. In another

experiment (not shown) the iron clusters grown at

150 K were annealed for 20 min at 350 K and

STM images were then taken at 150 K. Again,

the apparent cluster height was the same as that

attributed to 2D islands. Thus, the shape of the

iron clusters was determined by the sample

temperature during deposition and remained

unchanged during annealing to room temperature

or even slightly above. The information gained

here on the stability was an important

prerequisite to subsequently running catalytic

reactions or investigating the temperature

dependence of magnetic properties, such as the

zero field susceptibility [232]. The fact that the

growth and annealing temperature have quite

different influences on the cluster shape has been

previously observed in the case of Pd/MgO(100)

[115]. Such a phenomenon was caused by the

generally higher energy barrier required to

transform an entire island into its equilibrium

shape as compared to the one required to put the

atoms, being added one by one during growth

onto thermodynamically favorable sites, and

being lower coordinated. Therefore islands may

well remain 2D upon annealing to temperatures

higher than the growth temperature where they

are 3D. Figures 17 shows the bias voltage

dependence on the height and diameter of the

iron clusters [222]. The observed increase in

cluster height at around 1.5 V could be attributed

to empty states of the Fe clusters [233] whereas

empty states of the Al2O3 film contributed around

2.4 V [157].

4. Conclusions

Recent studies on the growth morphology of thin

films on the metallic and oxide surfaces have

been summarised. The investigation leads to the

following new results. In the first section of the

review growth behaviour of ultra-thin Ag, Au, Sn

and Pb films on Mo(110) was discussed as well

as, in addition to this, Pb on Ni3Al(111) surface.

For Ag on the Mo(110) surface, an analysis of

measurements indicated that three-dimensional

growth of a Ag film was observed. For

submonolayer coverage, the growth of Ag was

mediated by a two dimensional step-flow

mechanism. During the initial stage of this

growth, the first Ag layer nucleated and created

islands at Mo step edges. In the monolayer

coverage range, the decoration of substrate steps

by Ag could be distinguished by the presence of a

fractional step height at the Ag-Mo boundary. As

the sample was post-annealed to 700 K, the

morphology of the surface changed.The step-

flow growth in this case gave rise to a regular Ag

nanostripe network attached to Mo(110) step

edges. In case of Au and Sn layer-by-layer

growth for the first two layers was observed. For

submonolayer coverage, gold and tin created one-

atom high islands on the Mo terraces. Gold

nucleated preferentially along the ]100[ and ]111[

directions, however, tin did not exhibit a

preferential nucleation direction. In the case of

gold and tin no ordered regions were observed in

the completed first and second layer. As the

samples were post-annealed to 800 K,

rearrangement of existing Au and Sn films

occurred and Au-Mo and Sn-Mo surface alloys

were formed. Finally, for Pb films on Mo(110)

and Ni3Al(111) surfaces, the two-dimensional

growth of the first Pb monolayer wetting layer

took place below 1 ML. For > 1 ML, three-

dimensional growth of the Pb islands were

observed on the wetting layer with strongly

preferred atomic scale ‘magic heights’ and flat

tops. Here, at coverages between 1 and 2 ML,

STM showed the growth of two-atom high (N =

2) lead islands. Furthermore at higher coverages,

island height distribution showed peaks at

relative heights corresponding to N = (2, 4, 6, 7

and 9) and N = (3, 5, 7 and 11) atomic Pb layers

for Pb/Mo(110) and Pb/Ni3Al(111), respectively.

As the Pb/Mo(110) and Pb/Ni3Al(111) samples

were post-annealed to 1050 K and 400 K,

respectively, a change of the height of the lead

‘magic islands’ was observed. The island height

distribution exhibited peaks in the relative

heights corresponding to N = (2, 5 and 7) and N

= (2) atomic Pb layers above the wetting layer for

the Pb/Mo(110) and Pb/Ni3Al(111) respectively.

In the second section of this review the properties

of alumina films and the growth morphology of

Copper Phthalocyanine and Fe on the

Al2O3/Ni3Al(111) surface were focused on. The

structural relation between an ultra-thin Al2O3

film and the Ni3Al(111) substrate has been

investigated. It has been found that only the so-

called ‘dot structure’ is a real superstructure of

the Al2O3 film. The unit cell size of this

(6767)R47.784 superstructure (‘dot

structure’), bdot, has been determined with high

precision to be 41.6 Å. SPA-LEED as well as

LT-STM measurements clearly showed the

commensurability of the Al2O3 film and the

Ni3Al(111) surface. This revealed that the

structure of the alumina film is determined by


27

both the intrinsic structure of the film itself and

the structure of the substrate. In STM the

appearance of the Copper Phthalocyanine as

four-lobe protrusions was in good accordance to

the calculated shape of the lowest unoccupied

molecular orbital (LUMO). The weakly

chemisorbed CuPc molecules showed a high

lateral mobility above 250 K. At 350 K, thermal

desorption was complete. At very low coverages,

below 0.1 ML, a weak template effect of the

alumina surface was observed. Only two different

molecular adsorption orientations of CuPc were

distinguishable, which were rotated by 30 with

respect to each other. With increasing CuPc

coverage the strong intermolecular -electron

interactions compensated the weak template

effect of the surface resulting in the growth of 3-

DIM molecular clusters without any visible

ordering. The distance between the molecules in

the first and second layer was 3.5 Å, which is

similar to that in the -modification of bulk

CuPc. By the use of STS the degenerated LUMO

of the adsorbed CuPc molecules has been

identified at an energy of 1.2 eV above EF.

Finally, results presented here for growth

morphology of iron on the Al2O3/Ni3Al(111)

indicated that the nucleation site of iron clusters

was largely independent of deposition

temperature, while the cluster shape was 2D in

the range 130 – 160 K and 3D at 300 K. At

coverages close to one monolayer the Fe clusters

formed a regular hexagonal lattice with 24 Å

spacing, while at low coverages they occupied

only some of the available nucleation sites. In

comparison with Mn and V cluster growth on

alumina films, the observed perfection of the

order of the cluster arrays increases from Mn- to

Fe- to V-cluster arrays [106]. Iron clusters grown

between 130 K and 160 K were stable against

coarsening and shape changes upon prolonged

annealing to temperatures up to 350 K. Future

theoretical and experimental work is required to

further study the electronic structure and stability

of these uniform Pb ‘magic islands’ on Mo(110)

and Ni3Al(111) surfaces. This type of growth

could have impact on the engineering of stable

materials and devices with nanometer-scale

dimensions. Secondly, to confirm Sn-Mo and

Au-Mo alloy formation.

Acknowledgements

The projects described in this review were

financially supported by the Ministerium für

Wissenschaft und Forschung des Landes

Nordrhein-Westfalen (MWF-NRW), and the

Deutsche Forschungsgemeinschaft (DFG)

through SFB 624 and the priority program

“Organic Field Effect Transistors”, all of whom

are gratefully acknowledged. The work of A.

Krupski was supported by the University of

Wroclaw under Grant Nos. 2016/W/IFD/2006,

2016/W/IFD/2007, 2016/W/IFD/2008,

2342/W/IFD/10, 2016/W/IFD/10,

1010/S/IFD/10, 1010/S/IFD/11. The author also

gratefully acknowledges financial support gained

through a fellowship from the Alexander von

Humboldt and Hertie Foundations (III 2003–XII

2004, I-III 2011). A.K. permanent address:

Institute of Experimental Physics, University of

Wrocław, pl. Maxa Borna 9, 50-204 Wrocław,

Poland. The work reviewed here would have

been impossible without the contribution of my

colleagues. I am particularly grateful to Klaus

Wandelt for many insightful discussions and his

continuing support, which made part of this work

possible. I am furthermore greatly indebted to the

past members of his research group at Bonn

University who have contributed partly to the

results reviewed here, in particular Stefan Degen,

Marko Moors, Marko Kralj and Conrad Becker

for their contributions to studies of aluminum

oxide. I would like to thank all people who

contributed partly with experimental, theoretical

and technical assistance work for metal growth

on Mo(110), in particular Leszek Jurczyszyn,

Krzysztof Kośmider, Barbara Pieczyrak, Adam

Paruszewski, Wojciech Linhart, Dominika

Grodzińska, Izabela Cebula, Zbigniew

Jankowski, Mirosław Karpicki-Antczak, Jakub

Cichos, Sławomir Chomiak and his coworkers.

Finally, I would like also to thank to Katarzyna

Miśków for her contribution to the studies of Pb

on Ni3Al(111) and S. R. C. McMitchell for his

help and very useful discussions during

preparation this topical review.

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Growth morphology of thin films on the metallic and oxide ... › portal › files › ...metal films and body-centered cubic bcc metals is important in understanding the initial growth