Electronic and Geometric Properties
of Silver and Gold Nanoparticles
hv
e-
Tip
Dissertation
Zur Erlangung des akademischen Grades
des Doktors der Naturwissenschaften
An der Universtität Konstanz,
Mathematisch-Naturwissenschaftliche Sektion,
Fachbereich Physik
Vorgelegt von
Ignacio López Salido
Tag der mündlichen Prüfung: 29. Januar 2007
Referent: Prof. Dr. Gerd Ganteför
Referent: Prof. Dr. Paul Leiderer
Konstanzer Online-Publikations-System (KOPS)
URL: http://www.ub.uni-konstanz.de/kops/volltexte/2007/2707/
URN: http://nbn-resolving.de/urn:nbn:de:bsz:352-opus-27075
Cover picture: Nanoparticles on a substrate being investigated by means of
Spectroscopic and Microscopic Techniques.
List of Publications
1 Defect Formation of Au thin Films on SiO2/Si upon annealing:
D. C. Lim, I. Lopez-Salido, R. Dietsche and Y. D. Kim:
Philosophical Magazine, 2005, Vol. 85, N 29, 3477–3486.
2 Ag Nanoparticles on Highly Ordered Pyrolytic Graphite (HOPG) surfaces
studied using STM and XPS:
I. Lopez-Salido, D. C. Lim and Y. D. Kim:
Surface Science, 2005, Vol 588, Issues 1-3, 6-18.
3
Size Selectivity for CO-Oxidation of Ag Nanoparticles on Highly Ordered
Pyrolytic Graphite (HOPG):
D. C. Lim, I. Lopez-Salido and Y. D. Kim:
Surface Science, 2005, Vol 598, Issues 1-3, 96-103.
4
Oxidation of Au nanoparticles on HOPG using atomic Oxygen:
D. C. Lim, I. Lopez-Salido, R. Dietsche, M. Bubek, Y. D. Kim:
Surface Science, 2006, Vol 600, Issues 3, 507-513.
5
Electronic and Geometric Properties of Au Nanostructures on HOPG
studied using XPS and STM:
I. Lopez-Salido, D. C. Lim, R. Dietsche, N. Betram, and Y. D. Kim
Journal of Physical Chemistry B, 2006, Vol 110, Issue 3, 1128-1136.
6
Characterization of Ag Nanoparticles on Si wafer prepared using Tollen’s
Reagent and Acid-Etching:
D. C. Lim, I. Lopez-Salido, Y. D. Kim:
Accepted in Applied Surface Science (2006).
7
Size-Selectivity of the Oxidation Behaviors of Au Nanoparticles:
D. C. Lim, I. Lopez-Salido, R. Dietsche, M. Bubek, Y. D. Kim:
Angewandte Chemie International Edition, 2006, Vol 45, Issue 15, 24132415.
8
Experimental Studies on Plasmon Resonance of Ag Nanoparticles on
Highly Ordered Pyrolytic Graphite (HOPG):
I. Lopez-Salido, N. Bertram, D. C. Lim, G. Ganteför and Y. D. Kim:
Bull. Korean Chem. Soc. 2006, Vol 24, No 4, 6-562.
9
Electronic and Chemical Properties of supported Au Nanoparticles:
D. C. Lim, I. Lopez-Salido, R. Dietsche, M. Bubek, Y. D. Kim:
Submitted in Chemical Physics (2006).
Contents
Contents ...............................................................................................................1
1 Introduction......................................................................................................1
2 State of the Art..................................................................................................5
3 Scanning Tunneling Microscopy ..................................................................17
4 Photoelectron Spectroscopy ...........................................................................23
5 Experimental Setup........................................................................................31
5.1 Procedure.................................................................................................31
5.2 Equipment ...............................................................................................34
5.3 Supporting Material: HOPG and SiO2...................................................40
6 Results and Discussion ..................................................................................45
6.1 Ag and Au Nanoparticles on HOPG ......................................................45
STM Analysis ............................................................................................45
XPS Studies ...............................................................................................71
Summary....................................................................................................87
W-Oxide Coating on Ag Nanoparticles ....................................................88
6.2 Au Nanoparticles on SiO2/Si ..................................................................91
STM Analysis ............................................................................................91
XPS Results ...............................................................................................95
Summary....................................................................................................99
6.3 Bimetallic (Ag-Au) Nanoparticles grown on HOPG ...........................101
STM Analysis ..........................................................................................101
XPS Studies .............................................................................................107
Summary..................................................................................................109
7 Chemical Behaviours of Ag and Au Nanoparticles....................................111
8 Conclusion....................................................................................................115
9 Outlook .........................................................................................................117
10 Zusammenfassung .....................................................................................119
A List of Figures..............................................................................................123
B List of Presentations.................................................................................... 131
C References ................................................................................................... 133
Acknowledgment ............................................................................................. 143
1 Introduction
The creation and utilization of new materials, devices, and systems
through control of matter at nanometer scale (10-9 m), and the ability to work at
these levels with nanostructures is known as Nanotechnology. At least one
dimension at nanoscale is necessary to generate building blocks for these
nanostructures. Candidates for building blocks can be nanowires, thin films,
supramolecular aggregates, and nanoparticles. In this context, noble metal
particles consisting of two to about several thousand atoms have attracted
considerably the attention over the past few years. The reason is that these
particles at nanometer scale reveal novel properties different from individual
atoms, molecules and bulk materials. A basic understanding of these properties
opens fascinating routes to design devices and potential technological
applications. Some of these novel properties have been already explored and
exploited in photography [1], biological labelling [2], photonics [3],
information storage [4], optoelectronics [5], etc.
Gold as bulk, usually appreciated not only for its beauty but also for the
resistivity against corrosion, is a special case at nanoscale. In the form of small
nanoparticles, it becomes especially catalytic active at certain sizes. The
Industrial Technology Research Institute in Japan, for example, claims that gold
nanoparticles are ready to hit commercial markets for gas filters, air
conditioners, and air purifiers [6]. In other areas of applications like aerospace
technology, gold is used as protective coating of the Hubble telescope to reflect
heat flow whilst permitting light very well [7]. In electronics, for example,
single-electron switches and transistors of metal quantum dots1 are developed
with gold (or silver) nanoparticles.
In the case of silver nanoparticles, there are already many interesting
applications. The activity of silver as catalyst for the selective oxidation of
hydrocarbons, such as ethylene or methanol, is well known [8, 9], opening the
interest for silver nanoparticles as alternative material to the excellent properties
of gold nanoparticles in this field. In other fields, like pharmacology, silver is
employed as a bactericide for water purification and to prevent the buildup of
bacteria and algae in filters since more than a decade. Studies on the function of
1
Quantum dot: If we cut off a small piece of metal nanowire, we reach a zero dimensional
situation, where the remaining electrons are constrained in such a way that they begin to occupy
discrete energy levels. This particle can be only described by quantum mechanical rules.
1
2
1 Introduction
silver nanoparticles as antimicrobial agent are very promising [10]. Due to the
sensitivity of noble metals to light and temperatures of wide ranges, recently, in
the field of medicine, some attempts to use these properties for applying silver
(or gold) nanoparticles to fight against cancer, were pursued [11]. For example,
microcapsules filled with a drug, can be delivered directly to its target location
within a cancer cell, avoiding undesired side effects of medications. In order to
release the drug, these nanoparticles are introduced into the shells of the
capsules and the microcapsules can be heated in the presence of a fluorescence
dye. These metal nanoparticles act as absorption centers for this light so that a
laser pulse can be used to break apart the capsules and release their contents.
Along this line of medical applications, silver and gold nanoparticles in
combination with bio molecules are already at a point where even commercial
application in medical diagnostic has become known [12].
Many of these special properties are based upon quantum confinement
effects at nanometer scale. A basic explanation of this is the following: the
properties of a material depend on the type of motion its electrons can execute,
which depends on the space available for them. If the physical size of the
material is reduced to nanometer size, its properties change unpredictably,
becoming sensitive to its size, shape and electronic structure [13].
As already mentioned, one of the properties, which can vary extremely as
the size of the particle becomes very small, is its catalytic activity,. At
nanometer scale, this activity can depend on the geometric and electronic
structure of the particle. For supported nanoparticles, the influence of the
substrate should be also taken into account. The catalysis of noble metal
nanoparticles became especially important at the end of the 1980´s as M.
Haruta and co-workers discovered that gold nanoparticles can exhibit a high
catalytic activity for reactions such as CO-oxidation, propylene epoxidation and
other industrial reactions, specially at low temperature (about 200 K) [14-18].
This discovery caused excitement in the surface science community, opening
many questions for a better understanding of catalytic processes. It was
suggested that a characterisation of these particles could shed light onto these
catalytic behaviours. Since the substrate may reshape the geometry and
electronic structure of the particle, investigations of deposited particles on a
substrate are essential. In line with this idea, the results presented in this work
concentrate on the geometric and electronic characterization of silver and gold
nanoparticles supported on two different substrates.
3
1 Introduction
A promising method to investigate these properties is given via controlled
deposition of size selected clusters2. Clusters can be generated by cluster
sources (e.g. laser vaporization source [19], arc discharge source [20],
magnetron sputter source [21, 22]) in the gas phase, subsequently mass-selected
(via mass-spectrometer) and deposited on the substrate. Additional information
can be obtained, preparing relative narrow size distributions by nucleation and
growth on the support. Narrow size distributions of supported nanoparticles can
be achieved via conventional methods as metal vapour deposition [23], and
chemical methods such as synthesis of colloidal particles surrounded by organic
stabilizers [24] or by using the reduction of Tollen’s reagent and chemical
etching [25]. Thus, information gained from mass-selected clusters and
supported nanoparticles (not mass-selected but with narrow size distributions)
can be compared and complemented.
Finding out the mechanisms that make noble metal particles (e.g. gold
and silver) a very active catalyst, is still an open question. Many factors, which
can contribute to this activity, should be further investigated, for example, the
geometry of a metal nanoparticle with its atoms at edges or corners. It could
also occur that these particles have a special electronic structure or that the
charge of these clusters is relevant for the activity. Possibly, these nanoparticles
need to sit at special sites (such as a defect) on a specific substrate (e.g. oxide
support). In any case, the size of these particles seems to play a key role for the
catalytic activity and therefore, particles at nanometer scale should be prepared
on a substrate in order to be investigated. In this thesis, additional information
about these mechanisms for Ag and Au nanoparticles has been provided. To
investigate how these different factors are involved, two stable model systems
were prepared: Ag (or Au) nanoparticles on a Van-der-Waals surface and
Ag(Au) nanoparticles on an oxide substrate. Both systems were as well defined
as possible, resulting to be highly interesting as a potential catalyst for COoxidation. The characterization of the geometric and electronic properties of
these particles is an essential step to understand why these particles have a sizeselective catalytic activity.
Last but not least, it is worth mentioning that cleaner cars, non-polluting
industrial-process or eco-friendly power plants could be in a promising future
some of the wonderful and idealistic achievements based upon the knowledges
acquired by basic scientific works related to this dissertation. For this purpose,
2
Clusters are particles consisting of 1 to about several hundreds of atoms with properties that
diverge from single atoms, molecules or bulk features.
4
1 Introduction
it is still necessary to investigate these systems but at ambient conditions3.
In the next section, a short review over the last years is presented,
mentioning the most significant results on gold and silver nanoparticles and
mass-selected clusters. In doing so, I focus more on those aspects which are of
interest to this work and refer to the important results from the literature.
3
In this work, experiments were done under Ultra High Vacuum conditions (< 10-9 mbar).
2 State of the Art
After Haruta’s discovery [26], many studies have been focused on
explaining why gold is a good catalyst for the low temperature CO-oxidation4,
while Pt-metal catalysts, for example, need a higher temperature for this
oxidation reaction. Gold nanoparticles as model catalyst have been widely
investigated over the last years, experimentally as well as theoretically.
Although the highest number of studies is concentrated on gold, in principle,
the interest of the community is also recently attracted to Ag nanoparticles.
One of the first experiments demonstrating that the catalytic activity of
gold nanoparticles can be related to the electronic structure was performed by
Goodman et al.[15]. In this experiment, gold nanoparticles in a size regime
between 1 and 6 nanometers were prepared on a TiO2 (110) surface and
investigated using STM5. Goodman and coworkers found a maximum in the
catalytic activity for the CO-oxidation as the particle size becomes about 3-4
nm. Moreover, the catalytic activity was found to coincide with an onset of a
band gap (metal-nonmetal transition) in the electronic structure of the particle
(see Fig 2.1). Other groups have also predicted this kind of transition for these
nanoparticles on other supports. Metal-nonmetal transition was predicted with
X-Ray Photoelectron Spectroscopy (XPS, see Chapter 4) for gold nanoparticles
about 150 atoms by means of an extrapolation of the XPS data obtained by size
selected clusters (Au33, Au27, Au7 and Au5) supported on amorphous carbon
[27]. In this system, different shifts in the valence and core levels, depending on
the cluster size, were observed. The authors in ref [27] correlated these shifts
with the average coordination number of the atoms in the particle. In line with
ref [27], Whetten and co-workers found for gold nanoparticles supported on
Aluminia that discrete states (interband gap 5d-6sp), in the optical absorption
spectra began to emerge at diameter below 2 nm [28]. They noted that this
transition takes place as a step-like structure in the particle emerges. Atoms at
steps are undercoordinated, differing in the electronic structure from atoms in
the bulk. Therefore, it is reasonable that a reduction in the average coordination
number of the atoms partly affects the electronic structure of the particle [29].
4
5
1
CO + O2 → CO2
2
STM (Scanning Tunneling Microscope). For further explanations, see Chapter 3.
5
6
2 State of the Art
Theoretical studies about the dependence between the electronic structure of
undercoordinated atoms and geometry of the particle have been done by
U.Landman et al. [30]. They calculated that a strong hybridization of s and d
orbitals for Au cluster up to approx. 13 atoms is responsible for a planar
structure of these clusters.
50
40
30
nm
20
10
0
0
10
20
30
40
50
nm
Fig 2.1 Catalytic activity and electronic structure of gold nanoparticles on TiO2 (100) for the
CO-oxidation as function of particle size. By Goodman et. al.,[15]
Shifts of core levels in the particle to higher binding energies with
decreasing particle size have been frequently observed [31-37]. In addition,
oxidation states of the particles were often found to lead to core level shifts, or
new peaks at higher binding energies, related to new oxygen species of the
particle [38]. For these reasons, an experimental method, such as XPS, is a
useful method to study the electronic structure of the particle (see Chapter 4).
Initial and final states effects, before and after the photoemission of an electron
from the particle, have been suggested to explain these changes in the electronic
structure observed by XPS [39]. These effects are widely examined for Ag and
Au nanoparticles in Chapter 6.
In the cluster community, the possibility of creating mass-selected cluster
by means of diverse cluster sources [40, 41] or chemical methods has permitted
2. State of the Art
7
a systematic study on gradual changes in the electronic structure with
increasing cluster size on the atom-by-atom basis. It is well known, that a large
energy gap and a closed electron shell configuration is a prerequisite for the
chemical stability of a cluster. Thus, for example, the large HOMO-LUMO6
gap of C60 is responsible for its chemical inertness [42]. Au20 has also a closed
shell electronic structure [43, 44]. In case of Au55 (~1.4 nm) prepared on top of
a silicon wafer, a maxima resistance to be oxidized was found in oxidation
experiments performed by Boyen and coworkers [45]. In this work, this
resistance was not observed for gold nanoparticles with other sizes (between 1
to 8 nanometers), prepared in a similar way. In contrast to Au20, Au55 was
suggested to be magic due to geometric reason even though it is metallic in
electronic structure. For this reason, Au55 was proposed to be special candidates
as catalysts for the CO-oxidation.
Evidences in the gas phase, that the electronic structure can determine the
reactivity of Agn and Aun metal clusters were also reported by Ganteför et. al.,
[46-47]. In this work mass spectra taken for silver and gold cluster anions; i.e.,
Ag-n and Au-n (n=2, 3, 4… 19, 21) show a pronounced even-odd alternation of
O2 uptake, which is directly related to their electron affinity. This alternation is
shown in the case of free Au-n clusters in Fig 2.2.
-
Au3
-
Intensity (arb.units)
-
Au5
-
Au4O2
Aun + O2
-
-
Au6O2
Au7
-
-
Au3O2
3
4
-
Au9
Au8O2
5
6
7
8
9
Number of Au atoms
Fig 2.2 Mass spectra of free Au-n cluster anions for the mass regime (3 < n < 9) after reaction
with O2. Only cluster with even numbers of atoms reacted with O2. By D. Stolcic et. al., [46].
6
Energetic gap in a cluster between the Lowest Unoccupied Molecular Orbital (LUMO) and the
Highest Occupied Molecular Orbital (HOMO).
8
2 State of the Art
It is worth mentioning that time-resolved experiments on metal clusters in
the gas phase are becoming very important. Exciting studies about the dynamic
of the photodissociation of O2 molecules from Ag-n, and Au-n metal clusters
using spectroscopy methods such as Photoelectron Spectroscopy (PES), can
provide information in real time (femtosecond scale) about the underlying
mechanisms of O2 dissociation and CO-oxidation ongoing on these particles
[48-50].
Au3
Au2
Au6
Intensity /arb. units
Au5
88
Au7
Au9
Au8
92
Au4
84
80 9 2
88
Au10
84
80 92
B ind ing Ene rgy / eV
88
84
80
Fig 2.3 XPS spectra of Aun cluster (n=2 - 10) deposited on SiO2. An even-odd alternation in
the reactivity of this clusters for the absortion of atomic oxygen is observed . By D. C. Lim
et.al.,[51, 52].
The strong alternation of the chemical properties of the coinage metals
observed in the gas phase can become less pronounced, when clusters interact
with a substrate. Fig 2.3 shows XPS spectra measured by Y. D. Kim et al. [52],
in which this even-odd alternation is again observed. Low reactivity for Au3,
Au5 and Au7, and high reactivity for Au2, Au4, Au6 and Au8 is demonstrated.
Clusters with even number of atoms present additional peaks, at higher binding
energies respect to the Au 4f states, which are related to the new oxidation state
of the particle [51, 52]. Although this alternation still exists for all evennumbered clusters in the figure, attenuation in the alternation-pattern can also
2. State of the Art
9
be observed for bigger clusters, i.e. for n>8, the even-odd pattern is much less
pronounced than in the case of the respective gas-phase clusters. Scott L.
Anderson and coworkers discovered a strong size dependence on the CObinding and oxidation of Aun (n=1, 2, 3, 4, 7) clusters deposited on TiO2, which
is much different from the size dependence of the chemistry of gas phase
clusters [53]. In order to investigate the oxidation-pattern of larger particle size
regimes, Y.D. Kim and coworkers investigated supported Au nanoparticles
(between approx. 1 and 10 nm) on SiO2/Si, grown via evaporation [54]. A size
selectivity of the catalytic activity of these particles for the CO-oxidation was
found7. Only very small particles with the particle heights of 2-3 atomic layers
or less were found to be inert towards oxidation.
The fact that the chemistry of deposited clusters is much different from
the respective free clusters points towards importance of metal-support
interaction as well as electronic properties of clusters themselves. Campbell and
coworkers, for example, demonstrated the importance of the role of the
substrate by preparing Au particles with controlled thicknesses from one to
several monolayers on TiO2(110) [55]. In this work, the authors studied
oxygen-adsorption on Au/TiO2(110) as a function of Au particle thickness.
Experiments carried out with Thermo-Desorption Spectroscopy (TDS8) indicate
higher O2 desorption temperatures (741 K) for ultra thin gold particles than for
thicker particles (545 K). This implies that oxygen molecules bind much more
strongly to the thin gold particles and consequently oxygen molecules
dissociates more easily from these particles. Furthermore, it is discussed
whether the catalytic activity can be explained by the presence of active sites
(defects) on the support or by highly reactive sites on the particle (low
coordinated atoms). The catalyzed combustion of CO for Au-n (n≤20) clusters
deposited on MgO, using soft landing deposition, was reported by U. Heiz et
al.,[18, 56]. The chemical reactivity of these clusters was investigated using
TDS. In this work, a significant size-dependence of the CO-oxidation reactivity
of the Au-n cluster was found. Especially, Au8 and Au18 were found to be the
smallest and the highest catalytically active sizes, respectively. Size-dependent
CO-oxidation for Ag clusters has been also reported, elsewhere [57]. U.Heiz
and coworker concluded that electron transfer from the surface to the gold
cluster due to the F-center defects created by the vacancy of oxygen atoms can
7
For additional information , see Chapter 7
Thermo-Desorption Spectroscopy (TDS), also known as temperature programmed desorption
(TPD), is a UHV technique, which consists of observing desorbed molecules form a surface by
means of a mass spectrometer when the surface temperature is increased.
8
10
2 State of the Art
play an important role for the activation of the cluster as a catalyst. Au clusters
seem to be active for the CO-oxidation only if the clusters nucleate at oxygen
vacancies. The nucleation of Au clusters at oxygen vacancies suggests that
surface defects may alter the electronic configuration of the metal particle. This
effect enables the activation (or dissociation) of oxygen on the Au particles,
which increases the rate of the CO-oxidation reaction. Therefore, the role that
defects (i.e. oxygen vacancies) play for the activation of the reaction is very
important. Besenbacher and coworkers have demonstrated the correlation
between a decrease in the density of oxygen vacancies and the amount of
supported Au nanoparticles onto TiO2 (110) [58]. In this study, gold was
evaporated on a highly defective TiO2 (110) surface. STM images of the sample
were taken and DFT9 calculations were carried out in addition. These
calculations confirmed that the charging of Au cluster is very important for the
catalytic activity. This charging, in case of supported nanoparticles, can be
understood as a result of support-to-metal charge transfer. Besides U.Heiz
results, many evidences of the importance of the substrate for the activation of
CO-oxidation were reported in the literature. For example, large differences in
the activity between supported gold nanoparticles on Titania and Zirconia
(Au/TiO2 » Au/ZrO2) for identical particles sizes have been attributed to
different natures of the substrates, which can be explained due to possible
mechanisms, involving the adsorption of oxygen [59]. Moreover, it is important
to mention that the electronic structure of supported metal clusters can change
by introducing impurity dopant atoms into a cluster. U. Heiz and coworkers
used the system Aun/MgO, replacing one gold atom of an Au4 cluster by one Sr
atom. This exchange led to an enhancement in the catalyzed oxidation of CO
for the system Au3Sr/MgO [60].
As mentioned before, geometry effects can influence the electronic
structure of the particle. Shape effects were observed by Haruta´s investigations
showing a better catalytic activity of hemispherical gold particles as opposed to
more spherical particles using TEM10 investigation [61]. Moreover, the
surrounding area of the nanoclusters can be especially active for the activation
of oxygen. In the past, DFT calculations of O and O2 adsorption and CO
oxidation on gold surfaces have bee carried out by Mavrikakiss et al.,
9
Density Functional Theory (DFT) is a quantum mechanical method used to investigate the
electronic structure of many-body systems, in particular molecules.
10
Transmission Electron Microscopy (TEM) is an imaging technique whereby a beam of
electrons is focused onto a sample and an enlarged version of the sample appear on a
fluorescent screen.
2. State of the Art
11
demonstrating that these molecules prefer to chemisorb on stepped surface
Au(211) [62]. Fig 2.4 shows the calculated step density as a function of particle
size. A maximum in the step density is found as particle size is about 3 nm.
This may be correlated with the onset in the reactivity of small gold particles
predicted by Haruta and others [62. 63].
Fig 2.4 Calculated percent of edge Au atoms by particle as a function of particle size. By
Norskov et al., [62].
Fig 2.5 Interaction energy in eV per molecule for CO and O adsorption versus coordination
number for gold atoms in various geometries. By Jacobsen et. al., [64].
12
2 State of the Art
Another theoretical study carried out by Norskov and coworkers have
reported about calculations focused on the catalytic CO-oxidation of Au10
clusters depending on the geometry [17]. Au10 possesses low coordinated atoms
which are able to interact more strongly with adsorbates as shown in Fig 2.5.
Eventually, shape of the cluster can alter during a chemical reaction inducing
different isomers, which can have different energy barriers of activation. It is
possible that one of them can activate better the reaction [60].
Topographic investigation with a STM can be very useful for the
characterization of supported metal nanoparticles. Some studies have shown
that evaporated Au and Ag metal atoms, assuming to be monomer in the gas
phase, can suffer agglomeration upon reaching the surface, forming larger
clusters (»10 atoms). For this reason, a characterization of the size and
morphology of the particles after deposition is necessary. Burato’s group, for
example, investigated with STM cluster size distributions that result from the
deposition of Ag+n (n= 1, 2, 3) mass-selected clusters on TiO2 at room
temperature [65]. It turned out that Ag+n clusters sintered to form threedimensional islands of approximately 30 atoms in size for dimers and 50 atoms
for monomers.
Fig 2.6 Two STM images of Au+3 clusters on TiO2 (100)-(1x1) surface at room temperature:
(a) 140 Å2,; (b) 50 Å2; (c) Cluster size distribution. The bright spots are the clusters. The bright
stripes are the coordinated Ti atom rows separated by the bridging oxygen rows, which are
dark. The dim spots that appear on the bridging oxygen rows are bridging oxygen vacancies. In
this case, cluster agglomeration was not found. By Burato et. al., [66].
A similar experimental work was performed by the same group with Au+n
(n=2-8) mass-selected gold clusters on TiO2 [66]. STM pictures of soft landed
2. State of the Art
13
Au single atoms suggest aggregation of monomers into larger clusters. A nearly
uniform height was found for Au+n clusters with n bigger than one. The
preference of particles for certain adsorption sites was also observed. The
images suggest that these clusters remain intact after landing. Particle size
distribution and two STM pictures of deposited Au+3 clusters on TiO2 from
Burato’s work of ref [66] are shown in Fig 2.6.
In another work by Besenbecher and coworkers, Au atoms were
evaporated onto TiO2 (110) in different doses and temperatures, and the particle
size was determined using STM [58]. In absence of oxygen vacancies,
aggregation and sintering of gold metal particles to larger particles formed after
evaporation were found. There may be some differences between the processes
of depositing clusters and of evaporating atoms, which could justify different
behaviors among various studies. First, mass-selected clusters such as clusters
prepared by Burato and coworkers in ref [66] are charged. In addition, the
impact energy by the deposition of these clusters is about 1eV/atom, which is
above the thermal energy in the experiment of Besenbecher and coworkers of
ref [58]. The group of K.Kern studied the kinetic impact energy and substrate
temperature as a function of the cluster size [67]. For this, mass selected Agn
(n=1, 7, 19) were deposited onto a Pt(111) at low temperature. In this work the
kinetic energy per cluster atom was found to be the most important parameter
for a controlled deposition. The authors concluded that via energy dissipation
into a rare gas buffer layer, non-destructive deposition of soft landed clusters
can be achieved easier. This technique is widely used nowadays for the
deposition of clusters. Information about the experimental setup for a
nondestructive deposition of mass-selected clusters has been provided by
W.Eberhardt et.al., [68] and U.Heiz et. al., [69]. It is worth mentioning that the
stabilization of nanoparticles can be generated artificially through defects
created either by the impact process or by sputtering. This was reported by R.
E. Palmer et. al., [70]. In this work, Ag+n clusters (n=2700) were deposited on a
previously sputtered Highly Ordered Pyrolitic Graphite (HOPG) surface. A
stable and randomly dispersed array of Ag2700 cluster on the surface was
observed.
There is a high interest to study individually, that is, locally, nanoparticles
grown on surfaces. The morphology and electronic structure of individual noble
metal particles supported on surfaces can be well characterized with Scanning
Tunneling microscope (STM) and Scanning Tunneling Spectroscopy (STS) (see
Chapter 3). For example, STM images in situ of the nucleation and growth of
14
2 State of the Art
single Au nanoparticles supported on a reduced TiO2 substrate have been
provided by H. J. Freund et al., in the past [71]. In this work, a comparison,
cluster-by-cluster, was made of the morphological evolution and stability of
these particles during thermal annealing. Other experiments, such as studies on
the luminescent process at single clusters generated by tunneling electrons with
sufficient energy have been also carried out. For example, H. J. Freund et. al
performed STS experiments on individual Ag metal particles grown on
Al2O3/NiAl (110) for this aim [72]. In this work, light emitted from the Ag
particles (Mie plasmon resonance11) was stimulated by electron injection from
the tip of the STM. Moreover, STS can be employed to study the electronic
structure of single particles, such as its band structure. Thus, information about
quantized electronic states on the facets of large particles grown on a substrate
can be obtained by using a combination of UPS12/STS measurements. Using
this methods, Hövel et. al., investigated the electronic structure of the Fermi
level for Au nanoparticles deposited on a defective HOPG surface at low
temperatures [73]. UPS spectra were compared with STS data, providing
additional information about the Local Density of State of the electrons (LDOS)
in the Fermi level of the particle This comparison is very useful, because the
electronic structure of the tip usually perturb and complicate the interpretation
of STS spectra taken at single particles [74, 75]. Similar information about the
electronic structure of the surface of Ag nanoparticles grown on Ag(111) has
been reported by W. Schneider and coworkers by means of a low-temperature
STM [76]. All of these examples illustrate the capability of the STM to acquire
detailed geometric and electronic information from nanostructures (islands,
nanoparticles, mass-selected clusters) on surfaces, highlighting the potential of
this tool at local scale.
In this chapter, some of the accomplishments of several groups worldwide
to understand the size-dependent electronic, structural and chemical properties
of noble metal clusters on different substrates were summarized. It was shown
that information about the electronic structure of Ag and Au clusters in the gas
phase can be very valuable to understand further relevant results about some
aspects of the catalysis of these particles deposited on diverse substrates. A
considerable number of measurements on Ag and Au metal clusters have
revealed that among the effects responsible for these properties are metal to
nonmetal transitions as well as changes in geometry, core level binding energy
11
12
See Chapter 4
Ultraviolet Photoelectron Spectroscopy (UPS). See Chapter 4.
2. State of the Art
15
shifts, etc., each depending on particle size. Many models of supported noble
metal catalysts have been widely investigated using spectroscopic techniques
(eg. XPS, UPS) and microscopy studies (STM) over the last years. In addition,
many investigations on reactivity and spectroscopic/microscopic studies on
ultra thin oxide films result to be of special interest, yet the conductivity of thin
oxide films on metal substrate is sufficient to use STM. Size-dependent
catalytic behaviours of Au (or Ag) metal clusters have been frequently found,
giving rise to the interest of the cluster community on nanocatalysis. STM
experiments have achieved to follow the growth and sintering kinetics on a
cluster by cluster basis, providing already significant results. However, some
challenges lie ahead. It has been a purpose of this dissertation to better
understand the different role of oxide- and Van-der-Wals-supports, which can
significantly influence the electronic and geometric properties of Ag and Au
nanoparticles. A deeper insight into both properties is the key for unveiling the
size-selectivity in the reactivity of these particles toward atomic oxygen and
subsequently CO-reduction of metal nanoparticles.
16
3 Scanning Tunneling Microscopy
In 1986 Binnig and Rohrer were awarded with the Nobel Prize for the
invention of Scanning Tunnelling Microscopy (STM) [77, 78]. Two decades
later, STM has become one of the most important techniques in surface science.
Information about the topography and the electronic structure of a surface at
atomic scale can be achieved with this technique [79-84].
Control voltages for piezotube
Tunneling
current amplifier
Feedback system
Z
Y
X
Sample
Tunneling
voltage
Data processing
and display
Tip-surface
interaction
Fig 3.1 Main principle of STM. A feedback mechanism for the piezo controls the distance
between tip and sample. Tunneling current and voltage can be monitored as the tip scans over
the surface.
The basic principle of a STM is presented in Fig 3.1. A sharp conductive
tip is brought very close to the surface (0.5-2 nm). At this distance, a voltage
difference V (bias voltage <<4 V) is applied between the tip and the surface
(the sample). As a consequence of the tunneling effect (explained later), a
current of electrons with intensity I (0.01 nA-50 nA) flows through the vacuum
gap (potential barrier). A measurement of this current at constant height
(Constant Height Mode), or, of the voltage at constant current (Constant
Current Mode) can be done while the tip scans over the surface. To be able to
adjust the distance between tip and surface, the tip is attached to a piezo-electric
17
18
3 Scanning Tunneling Microscopy
element, which has the property of changing its length when an electric field is
applied. By adjusting the voltage on the piezo element, the distance between the
tip and the surface can be regulated. The combination of three piezo elements
enables a fine adjustment of the tip’s height over the sample in the X-, Y-, and
Z-directions.
The Tunneling Effect
The physical idea behind STM is based upon a quantum-mechanical
effect called tunneling effect. In classical physics, an electron cannot penetrate
through a potential barrier if its kinetic energy E is smaller than the potential
height Φ of the barrier. A quantum-mechanical treatment of this problem
predicts that the electron has a certain probability to traverse the barrier and
reappear on the other side. This effect is based on the wavelike nature of the
electrons. When an overlap between the electron wave functions of the tip and
substrate take places, there is a certain probability of finding an electron of the
sample in the tip, or vice versa.
In order to introduce this quantum-mechanical treatment, the Schrödinger
equation can be applied, in a simple case, to calculate the spatial wave function
of the electron with an energy E that collide with a rectangular barrier, with
height Φ, for the spatial interval (- ∞ <x<+ ∞ ) as shown in Fig 3.2. After
obtaining the wave functions, the probability that the electron will be found in
any of the regions, A, B, C, can be calculated by squaring the absolute value of
the wave function.
Fig 3.2 Rectangular barrier with a potential height Φ. From left to right, an electron collides
with the barrier with an energy E lower than Φ.
3 Scanning Tunneling Microscopy
19
It turns out that the solution of Schrödinger's equation in the barrier
region is of the form:
Ψ B ( x) ∝ Ψ ( 0 ) e − kx where k =
2m(Φ − E )
h
0< x<a
(1)
The current is proportional to the probability of electron tunneling
through the barrier;
I ∝ e −2 kx where k =
2m(Φ − E )
h
0< x<a
(2)
Thus, equation (2) predicts that the tunneling current decays
exponentially within the barrier and a finite probability of transmission can be
expected [85]. This result is surprising, considering that a flux of electrons
through the barrier is not allowed with classical arguments.
In a similar way, this tunneling effect is present at STM. An electron
located at the Fermi level13 needs an additional amount of energy, known as
work function Φ, for leaving the metal. When the tip is electrically connected to
the sample, the Fermi Levels of sample and tip are aligned, and the energy band
diagram of the STM tunnel-junction can be represented as shown in Fig 3.3 (a).
Fig 3.3 Energy band diagram of a STM tunnel junction (a) before applying a bias voltage at
equilibrium and (b) after applying a positive voltage relative to the sample. In this case, an
electron current flows from tip to sample.
13
The Fermi Level is defined as the highest occupied molecular orbital in the valence band at
0K.
20
3 Scanning Tunneling Microscopy
The slanted top of the potential barrier is, in this case, a consequence of
the different work functions of the two metals. If we apply a bias voltage to this
tunneling junction by applying a positive voltage, Vbias, onto the sample, a
tunneling current is established. This current occurs because electrons of the tip
are free to tunnel into the unoccupied states of the sample's conduction band. As
shown in Fig 3.3 (b), the electrons in the tip, which are responsible for the
current, have energies between the Fermi energies of sample and tip.
For the barrier in Fig 3.3 (b), an average work function can be defined:
Φ=
Φ sample +Φ tip
2
(3)
; where E=eVbias
In the case of a variable potential barrier, this approximation can be done
taking a succession of square barriers, and solving for each square in the same
way. This method is known as WKB approximation [86, 87].
Applying this average work function given by equation (3) in the eq. (2), results
in
I ∝ e −2 kx where k =
2m(Φ − eVbias )
h
0< x<a
(4)
One can estimate from equations (3) and (4), that changes in tip-sample
distance of about 0.1 nm induce current changes by one order of magnitude.
The vertical resolution depends solely upon the z dependence of the interaction
between tip and sample. Because of its exponential form, STM achieve a
vertical resolution, neglecting other factors, of about 1.0 x 10-2 nm [88, 89].
Equations (3) and (4) are based on the assumption that a single atom is
responsible for the STM junction contact.
Electron tunnelling can be also understood as an interaction between
occupied and unoccupied states. This idea was developed by Tersoff and
Hamann’s model (T-H model) [90, 91]. They show that the probability of an
electron in state ψ1 at energy E1 to tunnel into state ψ2 at energy E2 partially
depends on whether there is an unoccupied state with the same energy in the
other electrode. Thus, the tunneling current represents electron transfer from the
filled state of the tip to an empty state of the sample, or vice versa. Fig 3.4
shows the generation of tunneling current by applying a negative bias voltage
on to the sample. In this case, the occupied states of the sample generate the
current.
3 Scanning Tunneling Microscopy
21
Fig 3.4 Schematic energy diagram of electron tunneling with respect to the density of states of
the sample. The occupied states of the sample (indicated with dark color) generate the current
[92].
In the T-H model, the tunnelling current I(V) can be described as a
function of bias voltage V by
∞
I (V ) ∝ ∫ ρ s ( E ) ρt ( E − eV )Τ( E , V , z ) ( f ( E − eV , T ) − f ( E , T ) ) dE
−∞
(5)
where ρs(E) and ρt(E-eV) are the LDOS14 for the sample and the tip,
respectively. The factor T(E,V,z) is the transmission probability of the
tunnelling gap. f(E-V,T) and f(E,T) are the Fermi functions15 before and after
applying a bias voltage V.
Neglecting the influence of temperature for small variations and
approximating the Fermi functions by step functions, equation (5) simplifies to:
∞
I (V ) ∝ ∫ ρ s ( E ) ρt ( E − eV )Τ( E , V , z )dE
−∞
(6)
If the bias voltage V is close to the Fermi level (eV<<Φ), T can be
14
Local Density Of States (LDOS) is the distribution of number of electrons allowed per energy
level as function of which energy is considered.
15
The Fermi function, f(E, T), is the probability that a given available electron energy state will
be occupied at a given temperature.
22
3 Scanning Tunneling Microscopy
approximated by:
⎛
⎞
m
T ( z ) = exp ⎜⎜ −2 z 2 (Φ s + Φ t ) ⎟⎟
h
⎝
⎠
(7)
and the Intensity is given by
∞
I (V ) ∝ Τ( z ) ∫ ρ s ( E ) ρt ( E − eV )dE
−∞
(8)
It is evident from equation (8) and Fig 3.4 that information about the local
electronic structure of the sample can be obtained by measuring changes in the
intensity as function of the position on the sample. The changes in tip height
with position under feedback control reflect both the tip-sample separation and
the spatial variation of the local density of surface states (LDOS) of the sample.
Thus, the constant current image only reflects true height changes if the LDOS
of the surface (the local work function) is constant across the surface. This
would be the case for atomic steps on clean metal surfaces, but would NOT be
the case for adsorbates on surfaces. By altering the negative voltage applied to
the surface, other states contribute to the current. This provides a way to
calculate the LDOS of the sample. The first derivative dI/dV, which is
proportional to the LDOS, can be obtained from the measured I-V curves taken
at any spatial position. Thus, spectroscopic data corresponding to the
differentiation dI/dV normalized by I/V (i.e., (dI/dV)/(I/V)) contains
information about the electronic structure of the surface. This technique based
on STM measurements that gives information about the LDOS at atomic or
molecular scale is known as Scanning Tunneling Spectroscopy (STS). STS
measurements are usually carried out at lower temperature, because the spectral
resolution improves with lower temperature (∆E=3kBT) [93]. For example,
several phenomena related with magnetism [94, 95] and superconductivity [96,
97] could be traced by STS at temperatures around 4 K. A recent general
review on STS is given by Schneider et. al.,[93].
4 Photoelectron Spectroscopy
Photoelectron Spectroscopy (PES), or, Electron Spectroscopy for
Chemical Analysis (ESCA), is an experimental surface technique used to
determine the composition, the nature of chemical bonds and the electronic
structure of a surface region of a sample. It is based upon the photoelectric
effect, explained by A. Einstein in 1905 [98], in which monochromatic
electromagnetic radiation (photons) is used for the excitation of bound electrons
above the vacuum level. The small escape depth of these electrons makes this
technique surface-sensitive (see Fig 4.1).
Fig 4.1 "Universal curve" of the electron Inelastic Mean Free Path (IMFP=λ) versus kinetic
energies for different materials. By J.T. Yates et al., [90, 100].
Depending on the energy of the photon, usually in the range of X-rays
radiation (100 - 1500 eV) and ultraviolet radiation (5 - 40 eV), deep-core
electrons from the atoms of the sample or electrons from the valence levels can
be photoemitted, respectively. Using a synchrotron source [101], an entire
range of photon energies with high resolution enables more extensive energy
ranges of radiation (between 5 and 5000 eV). One of the advantages of working
with synchrotron radiation is the possibility of using an extremely narrow linewidth of the incident radiation. In contrast, the resolution for traditional X-ray
sources can be only improved by using X-ray monochromators.
23
24
4 Photoelectron Spectroscopy
In X-ray Photoelectron Spectroscopy (XPS), an electron in a core level
absorbs a photon with energy higher than its binding energy, and is emitted.
The kinetic energy of the photoelectron EK is related to the energy of the photon
hν by the following expression:
EK = hν − EB
(1)
where EB is the electron binding energy in eV, h is the Planck constant, and υ is
the frequency (Hz) of the radiation.
It is important to mention that the binding energy EB is referenced to the
Fermi level. The additional amount of energy that an electron needs to
overcome the energy difference between the Fermi energy (EF) and the vacuum
energy (Evacuum) is known as work function of the sample (Φsample). Considering
this additional energy, the kinetic energy (Ek) that an electron has after leaving
the sample, is given by
EK = hν − EB − Φ sample
(2)
Fig 4.2 Schematic energy diagram of X-ray photoelectron spectra. The measured kinetic
energy of the ejected electron is given by E*k. However, the true kinetic energy of the electron
leaving the sample is Ek.
4 Photoelectron Spectroscopy
25
Since the spectrometer possesses its own work function (Φspectrometer), Ek is
not the kinetic energy, which is measured by the spectrometer. If the sample is
electrically connected to the spectrometer, their Fermi energies are at the same
level as shown in Fig 4.2. The kinetic energy Ek* measured by the spectrometer
is given by:
Ek* = hν − EB − Φ spectrometer
(3)
It is important to mention that Φspectrometer can be measured with a bulk sample.
One-electron approximation
The first approximation to explain the photoemission effect of a single
electron was given by Koopmans Theorem [102] (also called One-electron
approximation16). In this approximation, the binding energy of an electron is
considered to be equal to the energy difference between initial state (atom with
n electrons) and final state of the atom (atom with n-1 electrons (ion) and free
photoelectron). Therefore, the binding energy of this electron is equal to the
negative value of the orbital energy ε:
EB = EFinal (n − 1) − EInitial (n) = −ε
(4)
In this approximation, no relaxation effects of surrounding electrons due
to the positive hole created by the ejected photoelectron are considered. This
simple picture of the photoemission effect implies, that rearrangement of all the
electrons around is neglected, considering as they were frozen. In addition, this
approximation neglects relativistic effects and effects of electron correlation.
Initial state effects
When the energies of core levels are investigated in detail; small shifts in
the binding energies of the electrons can be found. The initial state structure of
a particle17 can induce core level shifts. For example, chemical bonding
16
It is based on the Hartree-fock approximation. This quantum-mechanical approximation is
used to calculate the wave function of an electron orbital, assuming that the rest of the electrons
of the system are “frozen”.
17
A particle can be atom, cluster or molecule and can be free or supported on a surface.
26
4 Photoelectron Spectroscopy
influences the electronic configuration in and around the atom. In addition,
oxidation state and electro-negativity of neighboring atoms can have also
influence the electronic configuration, generally resulting in the appearance of
shifts or shoulders at the main peaks [103, 104]. Hence, direct information from
these shifts provides a quantitative method of chemical analysis. The ability to
resolve atoms exhibiting slightly different chemical shifts is limited by the peak
widths. Apart from the instrumental resolution (see Chapter 5), the width of the
photoemission peaks is determined by the lifetime of the positive core hole
created by the photoemission process. By the Heisenberg uncertainty relation,
the intrinsic peak width Г, is inversely related to the core hole lifetime τ by:
Γ=
h
τ
(5)
where h is the Planck constant.
As consequence of the statistical nature of the screening process of the
positive hole, the deeper the core hole, the more de-excitation channels there
are which can fill the core hole. This intrinsic lifetime broadening is usually
assumed to be of Lorentzian nature.
Final state effects
Final state effects reflect the relaxation energy of the system, which
corresponds to the energy difference between the excited electron system of a
particle (atom, cluster, molecule, or bulk) after losing a photoelectron and the
relaxation of the electron system. There are some differences depending on the
material. In a metal solid, i.e. metal bulk, the hole state (positive charge state)
created by the photoemitted electron, is completely shielded by the conduction
electrons and the cores of neighboring atoms. For isolated particles, in
particular, for deposited particles on poorly conducting substrate [34, 35, 105,
106], the number of conduction electrons and neighboring atoms is limited by
the particle size and, hence, the positive hole can be screened less efficiently.
Inelastic background
In a typical XPS spectrum, binding energy is plotted versus intensity of
photoemitted electrons. Fig 4.3 shows a XPS spectrum taken over a relatively
wide range (0 -1400 eV), obtained from an Ag bulk sample using Al-Kα
4 Photoelectron Spectroscopy
27
Ag (4d)
Ag (4p)
Ag (4s)
Ag (3s)
Ag (3p)
Ag (Auger)
Electron counts
Ag (3d)
radiation of 1486.6 eV. As Fig 4.3 depicts, the peaks ride on a background of
secondary electrons arising from higher kinetic-energy electrons. As result, the
background increases step-like beyond each major peak. This feature is due to
inelastic electron energy loss that happens as electrons from deep core levels
with a depth over the IMFP loose their kinetic energy. This background signal
can be subtracted by smooth curve-fitting, but may introduce some error in
quantitative intensity determinations and peak positions [107]. Moreover, all
peaks exhibit inelastic tails toward higher binding energies [108].
Binding Energy (eV)
Fig 4.3 XPS spectrum obtained from a Ag bulk sample using Al-Kα radiation [108, 109].
Auger Peaks
At higher binding energy in Fig 4.3, other peaks known as Auger peaks
appear. They are associated to electron transitions between core and valence
electrons and consequently to the emission of low energy electrons in an Auger
process. This process relies on the coupling between electrons in different
energetic levels, in contrast to the one-electron picture described before. An
Auger process is initiated by creation of a core hole, for example in the Kshell18, from where a phoelectron (Auger electron) is emitted above the vacuum
18
K, L, M are the quantum numbers n= 1, 2, 3,… respectively. Inside a shell (n=1, 2, 3,…),
levels (p, d, f,…) with a non-zero value of the orbital angular momentum ( l > 0), show spinorbit splitting and therefore two different energetic states (2p (L2,3), …, 3d (M2,3)…).
28
4 Photoelectron Spectroscopy
level. The ionized atom is in a highly excited state and tends to relax back
rapidly (ca. 10-14s) to a lower energy state. Thus, one electron falls from a
higher level L1 to fill an initial core hole in the K-shell and the energy liberated
in this process is simultaneously transferred to a second electron in the L2,3shell. A fraction of this energy is required to overcome the binding energy of
this second electron; the remainder is retained by this emitted Auger electron as
kinetic energy. In this Auger process illustrated in Fig 4.4, the final state is a
doubly-ionized atom with core holes in the L1 and L2,3 shells. The notation for
the transition illustrated in Fig 4.4 is KL1L2, 3 and can be energetically
expressed by:
E KL 1 L2,3 = E K - E L1 - E L2,3
(6)
Fig 4.4 Energetic scheme of the Auger process KL1L2,3 carried out using X-ray radiation.
(Left) The initial state (middle) ionization process occurs by removal of a K-shell electron. One
electron falls from a higher level (L1) to the core hole in the K-shell and the energy liberated in
this process is simultaneously transferred to a second electron in the level L2,3 (Auger electron).
(Right) Final state is a doubly-ionized atom [110].
Other Auger transitions for the scheme of Fig 4.4, are also possible, like for
example, KL1L1, KL2,3L2,3, L1L2,3L2,3. In general the analysis of the Auger
peaks can be used to detect elements. It is worth mentioning that Auger peaks
always accompany XPS, but they usually rise with broader and more complex
4 Photoelectron Spectroscopy
29
structure than photoemission peaks. Moreover, the kinetic energy of the Auger
electrons is independent of the incident energy of the radiation.
Spin-orbit splitting
The electronic structure of an atom can be described considering the
orbital and spin momenta of its electrons. It is well known that, for any electron
in an orbital with orbital angular momentum, a coupling between magnetic
fields of spin s and angular momentum L exists. For example, an electron is
removed from the 4f-level of an Au atom; the remaining 4f electrons may have
a spin either parallel or antiparallel to that of the remaining unpaired 4felectron. Because of different energy configurations in the coupling of two
parallel spins (j=L+s=7/2) and two antiparallel (j=L-s=5/2), these
configurations give rise to two states, resulting in a peak splitting. This splitting
can be observed in the Au 4f XPS spectrum (see Chapter 6). In this case, Au
4f7/2 states are at lower binding energy than the state Au4f5/2. For other orbitals
like the s orbitals, no angular momentum exists; therefore s orbitals do not show
spin-orbit splitting. These states are called singlet in XPS. In contrast, p, d, f
orbitals with angular momentum of 1, 2, 3, show spin-orbit splitting. These
states are called doublets.
Satellite peaks
X-Ray emission is generated by electronic transitions inside a metal.
Depending on the material of the X-ray source, different emission lines are
generated. These transitions, in which a photon is emitted, provide a
characteristic radiation at fixed photon energies. In addition, a continuous
background radiation of lower intensity known as “Bremsstrahlung” is
observed, too. Monochromatic radiation is very important to generate sharp
photoemission lines in XPS. The most popular monochromatic radiations are
due to 2p3/2→1s and 2p1/2→1s transitions, which originate from Al and Mg as
X-ray source, providing photon energies of Al-Kα1,2 (hν=1486.6 eV) and MgKα1,2 (hν=1253.6 eV). The same transitions in doubly ionized Mg or Al
generate Kα3,4 radiation energies of about 9-10 eV higher, inducing “satellite
peaks" in XPS spectra at lower binding energies. These satellite peaks are
shown in Fig 4.5 for a XPS spectrum of the C1s state of a HOPG surface.
30
4 Photoelectron Spectroscopy
Shake up, shake-off
In these events, the outgoing photoelectron excites a valence electron to a
previously unoccupied state, or, especially in a metal, to an electron-hole
formation (shake-up). An excitation of the valence electron above the vacuum
level is also possible (shake-off). For these transitions, the photoelectron must
give up some of its kinetic energy; hence, new features in the XPS spectrum
always lie on the high binding energy side of a direct photoemission transition.
Sometimes, these features don’t have the shape of discrete peaks, because
photoelectrons tend to fall into the energy region of inelastic secondary
electrons and often show no discrete structure. Fig 4.5 shows this feature in an
XPS spectrum of C1s state of a HOPG surface.
Plasmon
Photoelectrons may give up some energy to the electron gas situated in a
valence band of a conductor before leaving the material. This energy can be
transferred to the electron gas in form of collective oscillations known as
Plasmons, with a characteristic frequency. Due to this oscillation, peaks with
higher binding energies than the original binding energy can be observed in the
photoelectron spectrum (see Fig 4.5).
HOPG C1s
Intensity (a.u.)
Al Kα1,2
hv=1486,6 eV
energy loss
plasmon
345
330
315
Shake-up
300
285
Satellite
(Kα,3, Kα,4)
270
Binding energy (eV)
Fig 4.5 XPS spectrum of the C1s state of a HOPG surface. Satellite peaks at lower binding
energies are observed. In addition, Shakeup, shakeoff effects and plasmon oscillation can be
observed at higher energies [22, 111].
5 Experimental Setup
5.1 Procedure
All experiments were performed under Ultra High Vacuum (≤ 1 x 10-9
mbar) conditions at room temperature. Two UHV systems, (1) and (2), were
used for the STM and XPS experiments, respectively. System (1) is equipped
with a room-temperature STM (OMICRON STM1) located in the main
chamber. In the preparation chamber a Cylindrical Mirror Analyzer (CMA) for
Auger Electron Spectroscopy (AES) measurements, a He UV Lamp, a Low
Energy Electron Diffraction system, two electron beam evaporators (Tectra)
and a sputter-gun are situated. The sputter-gun is installed to clean HOPG
samples and create defect sites. In this work, 0.5 kV were used to accelerate Ar
ions towards the surface. The defect density was controlled by altering the
sputtering time and the Ar pressure in the chamber. The sample current during
sputtering was typically 1 µA, while Ar pressure of 2 x 10-5 mbar was used. The
estimation of the defect density created by sputtering, under these conditions, is
less than 0.5 % of a monolayer per second. This estimation relies on the
estimation of the attenuation in the intensity of the HOPG surface plasmon
resonance (C1s state) observed after sputtering the HOPG sample by means of
XPS. The XPS spectrum of C1s state of a non-sputtered HOPG surface is
shown in Fig 4.5. The HOPG samples were prepared by the scotch-tape pilling
and inserted into a sample-holder (home-made), located inside the preparation
chamber. Here, the sample can be outgassed at about 1100 K (see Fig 5.1). In
contrast, polycrystalline Si wafer samples with native oxide layers were
outgassed at about 800 K. Temperatures were measured with an optical
pyrometer (Impact). The cleanness of HOPG was confirmed using the STM in
UHV system (1) and XPS in system (2) (X-ray source: Al Kα; photon energy =
1486.6 eV). In the case of Si wafer, the cleanness was confirmed only by XPS.
For the preparation of a sample, two electron beam evaporators were installed
in the preparation chamber of the UHV system (1) (see Fig 5.1). Using both
evaporators, Ag (or Au) nanoparticles were grown on the sample by
evaporating Ag (or Au) atoms from an Ag (or Au) rod (purity 99.999 % from
Alfa Aesar) by means of electron bombardment heating19.
19
Electrons from a hot tungsten filament (thermionic emitter) are attracted towards a target
which is at high positive potential. A crucible contains the target which is evaporated.
31
32
5 Experimental Setup
The flux of Ag (or Au) can be kept constant by controlling the emission
current between W filament and Ag (or Au) target. It is worth mentioning that
from sample to sample, the emission current was varied between 15-23 mA to
alter the evaporation rate. Pt/Ir (90%/10%) tips were characterized with light
microscopy (Zeiss) and SEM (see Fig 5.2).
Fig 5.1 Sample-holder used in the preparation chamber of UHV system (1). In order to
tolerate high temperatures during outgassing process (up to approx. 1200K), the materials used
for this holder are tantalum and ceramic. Electrons from a hot tungsten filament (4-5 A) are
emitted and attracted towards a sample, to which a high positive potential (1-3 kV) is applied.
In addition, two electron beam evaporators are orientated directly to the sample, which can be
rotated and moved up and down.
250 nm
30 nm
0.15mm
Fig 5.2 Two images of a Pt/Ir tip taken with Left) Light microscopy and Right) SEM. By Ballot
[112].
5. Experimental Setup
33
After STM measurements, a small UHV transfer-system20 was used to
transfer the sample from UHV system (1) to UHV system (2). There, some
samples were further characterized.
In the XPS chamber, HOPG samples can also be cleaned. In this case, the
samples were outgassed at the temperature of about 800 K for more than 12
hours. The sputtering condition for this chamber is quite similar to those in the
preparation chamber in UHV system (1).
In the preparation chamber of system (2), Au nanoparticles were also
grown on the sample by evaporating Au atoms. In this case, a small Au piece
wrapped by a W filament, which can be resistively heated, was partly
evaporated. W-oxide coating was created by heating a W filament under an O2
partial pressure of 1 x 10-5 mbar. The preparation chamber has a base pressure
of 2 x 10-9 mbar. To provide atomic oxygen environments, samples were
exposed to a hot Pt-filament located at the backside of the sample in the
molecular oxygen atmosphere. Typically, Au nanoparticles were exposed to O2
pressure of 8 x 10-5 mbar for 30 minutes with a Pt-wire temperature of ~800 K.
The CO exposure experiments were also conducted in the preparation
chamber. After the gas exposure, the samples were immediately transferred to
the XPS analysis chamber, maintaining the pressure below 1 x 10-9 mbar while
the sample is transferred.
Fig 5.3 shows a scheme of the experimental procedure. In UHV System
(1), the sample can be prepared, and subsequently investigated in STM
chamber. A small UHV chamber is employed to transfer the sample between
the preparation chambers of both UHV systems. Inside a UHV system, the
sample can be carried to different positions in different chambers by means of
UHV manipulators.
20
See the UHV- Chamber inside the marked circle in the 3-dimensional drawing of Fig 5.4
34
Fig 5.3
5 Experimental Setup
Scheme of the experimental procedure in UHV System (1) and (2).
5.2 Equipment
Fig 5.4 shows a schematic 3 dimensional drawing of the UHV system (2),
in which STM investigations were performed. In the STM chamber the pressure
is maintained at about 2-4 x 10-10 mbar by an ion getter pump and a titan
sublimation pump. These pumps are installed at the preparation chamber. STM
5. Experimental Setup
35
and preparation chamber are directly connected with each other. The
preparation chamber is connected via a valve 1 to another chamber, in which
the sample can be inserted. This Chamber is pumped by a turbo molecular
pump. By means of a manipulator, the UHV transfer system can take the
sample from this chamber to its transfer chamber. After closing valve 3, the
transfer system can be separated and taken to the preparation chamber of UHV
system 2. During the time needed for transferring, the pressure in the transfer
chamber is maintained at 1 x 10-9 mbar by an ion getter pump.
STM Stage
Valve 1
STM Chamber
Manipulator
Manipulator
Valve 2
Transfer System
Preparation Chamber
Valve 3
Fig 5.4 Three dimensional drawing of the UHV system (2). Note: UHV transfer-system (see
UHV- chamber inside the marked circle are) was used to transfer the sample from UHV system
(1) to UHV system (2).
STM
From top of the STM chamber, the STM stage can be observed (see Fig
5.4 and Fig 5.5). The procedure to initiate an experiment consists of positioning
the sample and subsequently adjusting the scanner setup (software: Scala
Omicron Scala Pro). Scan mode, voltage gap, current, scan speed, scan range,
etc, are some of the parameters to be taken into account. In the next part, the
most important components of the STM equipment are described.
36
5 Experimental Setup
Surrounding column
of spring suspension
Permanent magnet
sample
Ring of copper
plates
Sample
slider
Piezos
tip
Fig 5.5 Picture of the STM stage taken from top view. Main components of the Omicron STM1
are indicated.
Spring suspension
The STM stage is suspended by four soft springs, which are protected by
surrounding columns. The resonance frequency of the springs is about 2 Hz.
The vibrations of the STM stage are intercepted using a non-periodic current
damping mechanism. This is achieved by surrounding the STM stage by a ring
of copper plates which come down between (permanent) magnets. The induced
current damps the vibrations.
Sample coarse positioning device
A CCD camera is mounted at the top viewport UHV flange of the STM
chamber. The sample can be coarse positioned close to the tip with help of this
camera. The sample slider is magnetically coupled to three shear piezos which
are driven with a voltage input. A slip-stick inertia effect is used here. The
sample slider is transported during slow movement of the piezo and slips during
fast piezo motion due to its inertia. With this technique, parallel and
perpendicular coarse movement of the sample stage to the tip can be achieved.
5. Experimental Setup
37
Wobble stick and carrousel
Tip- and sample-holder can be handled within the UHV STM chamber
using a wobble stick and carrousel. A carrousel is used to load (or unload) tipand sample-holders in UHV to transfer to and from the STM stage with the
wobble stick.
STM Tips
It is important to consider effects of the tips used for STM imaging (for
more information, see Chapter 3). For example, particle diameters are subjected
to the particle-tip convolution21. This results in a loss of accurateness on lateral
particle size. The schematic drawing in Fig. 5.6 (left) illustrates this. The bold
solid line on the left hand side represents an STM tip during scanning with a
tunnelling junction formed by an atomically protrusion, marked as a triangle.
250 nm
50 nm
Fig 5.6 Left: Schematic draw of a STM Tip scanning over a nanoparticle [113]. Right: SEM
image of the doubles apex of a bad etching W/tip. By H.Ballot [112]
On a flat surface, atomic resolution can be easier achieved, because the
tunnelling current is focused on the small “triangle”. But if a cluster of the size
of some nanometers is imaged, here indicated as a truncated sphere, the
tunnelling junction will take place somewhere else on the side of the tip contour
(see SEM image on Fig 5.6). This tunnelling junction takes place over the entire
tip while the cluster is scanned. In Fig 5.6, on the left side, the scanning process
is shown and the momentary position of the tip is indicated. The convolution
has two main effects. First, atomic resolution is lost in most cases. In the
21
Particle-tip convolution: Effect of the tip shape in STM imaging.
38
5 Experimental Setup
literature, there is only some exception as for example, in the work by H. J.
Freund and Besenbacher, et al., [114], where Pd nanoparticles with a diameter
of about 4 nm on a thin Al2O3 film on NiAl (110) were imaged with atomic
resolution. Secondly, a comparison of the tip trajectory and the real cluster size
results in an inflation of the lateral cluster size. This effect can be corrected by a
combination of XPS/STM measurements. This overestimation in the lateral
particle size was considered to estimate the particle size in this work (see
Chapter 6.1).
XPS
Fig 5.7 shows a schematic diagram of the UHV chamber, where XPS
measurements were performed. The sample is transferred from the preparation
chamber in XPS chamber and positioned perpendicular to the spectrometer with
an angle of 45° with respect to the light source (X-ray tube).
Fig 5.7 Schematic diagram of the experimental equipment in the UHV-XPS chamber. By
M.Bubek [51].
X-ray tube
From the X-ray tube (see Fig 5.7), a characteristic monochromatic
radiation (Al-Kα1,2 hυ=1486.6 eV) is emitted. This radiation is generated by a
high energetic electron beam of several keV, which strikes an Aluminium
anode. The electrons of the Al atoms get excited by the electron beam and L
electrons relax into the K shell, emitting photons with a characteristic
5. Experimental Setup
39
wavelength, known as Kα. A secondary, non monochromatic radiation, known
as Bremsstrahlung is emitted. The width of the line for the exciting X-ray
radiation Kα, is about 0.85 eV in this case. This value is very good compared to
the resolution of other X-ray anodes [115].
Electron energy analyzer
The energy distribution of electrons emitted from the surface can be
measured by an electron spectrometer. In this work, a Concentric Hemisphere
analyzer (CHA) from the company Omicron (EA125 U5) was used.
Schematic diagram of a Concentric Hemispherical Analyzer (CHA) for the detection of
the photoelectron energy in XPS [116].
Fig 5.8
Fig 5.8 shows a schematic cross section of the CHA. It consists of two
concentric hemispheres of mean radius R, mounted with common centre, to
which two different negative potentials are applied. The outer surface is more
negative than the inner, so that a mean equi-potential surface at R0 between the
hemispheres is formed. There are two slits located at the entrance and exit of
the hemispheres and both are centred on R. A lens system is located before the
slit at the entrance. These lenses are used to focus the emitting surface area. In
addition, the electrons get retarded before entering through the entrance slit into
the analyzer by an amount of energy Epass, which is kept constant, while the
40
5 Experimental Setup
electrons pass through the analyzer22. After passing through the exit slit of the
analyzer, the electrons are stopped by the detector. The resolution ∆E of the
CHA is expressed by
⎛r
⎞
∆E = E pass ⎜ + α 2 ⎟
⎝R
⎠
where r is the slit width and α the half angle for the electrons entering the
analyzer at the entrance slit.
The energy resolution is inversely proportional to the kinetic energy Epass
and to the radius R. However the sensitivity is directly proportional to the
intensity which decreases with increasing R and with lower kinetic energy.
Thus, a high energy resolution can be obtained at the expense of low sensitivity
(intensity). For this reason, a compromise in resolution and sensitivity is taken.
The pass energy Epass used for our experiments is 20 eV. The radius R is 125
mm, α is approx. 1° and the slit width accounts to 6 mm. Therefore, the energy
resolution of the analyzer CHA for our spectrometer in an energy range
between 100 and 1500 eV accounts to 950 meV.
5.3 Supporting Material: HOPG and SiO2
HOPG and SiO2/Si substrates were studied by a combination of XPS and
STM measurements in this thesis [103, 117]. The properties of these substrates
are summarized first.
Highly Ordered Pyrolytic Graphite, HOPG, possesses a layered structure,
with each layer composed of a planar arrangement of connected hexagonal
benzene rings. The triagonal bonding in each HOPG sheet involves overlap of
C sp2 hybrid orbital. Within a layer, HOPG has conducting properties. Amongst
C-atoms from different layers Van-der-Waals forces exist (See Fig 5.9).
22
This operation modus of the CHA, in which the analyser pass energy is held constant is
known as Constant Access Energy modus (CAE modus). In this modus the energy resolution is
the same for all possible kinetic energy values.
5. Experimental Setup
41
a
b
c
d
(a) Schematic diagram of the lattice structure of a small piece of HOPG with 3 atomiclayers. (b) STM image with atomic resolution of HOPG (2.2 x 1.1 nm2); STM parameter: 0.4V,
0.7nA. (c, d) 2- and 3-D images of a non-sputtered HOPG surface (243.4 x 243.4 nm2); STM
parameter: 0.4V, 0.5 nA.
Fig 5.9
Another important property of a HOPG surface is its surface plasmon23 at
6.4 eV, which can be observed by some surface techniques as HREELS (High
Resolution Electron Energy Loss Spectroscopy). This technique is based upon
the analysis of the energy spectrum of low-energy electrons scattered from a
sample under UHV conditions. Fig 5.10 shows a HREELS spectrum of a
HOPG surface, in which the π-plasmon resonance is observed. Using HOPG as
substrate for the study on growth and characterization of metal nanoparticles
has two main advantages. First, HOPG reacts very weakly with O2, CO, H2O
molecules in comparison with other substrates under UHV conditions; and
therefore sample cleaning is much less time-consuming compared to metal
oxide singles crystals or thin films preparations as SiO2, Al2O3, TiO2, MgO,
which are less inert materials [63]. Because there is almost no background
oxygen signal on clean HOPG surface by XPS and HREELS measurements,
23
Collective oscillations of electrons in the valence band of a conductor.
42
5 Experimental Setup
oxygen chemisorptions on metal nanoparticles supported on HOPG can be
well-studied [54, 103, 118, 119]. Second, metallic properties of HOPG allow
STM and XPS experiments without charging problems.
Intensity (a.u.)
HOPG
π-Plasmon
∆Ε
2000 x
-1
0
1
2
3
4
5
6
7
8
9
10
electron energy loss [eV]
Fig 5.10 HREELS Spectra of a non-sputtered HOPG surface; Lower curve: complete spectrum; upper
curve: enlarged electron energy loss feature- π- plasmon of graphite at ~ 6,4 eV. Incident electron-beam
energy E0 = 30 eV, specular impact angle θ=55°.
Fig 5.11 Two different STM images of the same sample of Ag nanoparticles deposited on nonsputtered HOPG. Left (496.2 x 462.5 nm2), Right (495.0 x 448.7); Tunneling parameter: 0.1 V,
0.2 nA. Note: Left and right images were taken successively with a time interval of 29 min
between both images.
In contrast to these advantages, it is difficult to acquire reliable STM
5. Experimental Setup
43
images of metal nanoparticles on HOPG because of a very weak metal-support
interaction [113]. Thus, STM tip induced movements of the deposited particles,
i.e. sintering or displacement of particles occurs on a Van-der-Waals surface
(See Fig 5.11). To avoid these effects, HOPG surfaces were sputtered in the
presented work.
SiO2 native oxide layer, i.e. SiO2/Si, were also used as supporting
material for metal particles to find out more about the role of different supports
and about the origin of any differences in the electronic structure of these
nanoparticles when metal atoms are evaporated on SiO2. In contrast to the
HREELS spectra of HOPG of Fig 5.10, in which no impurities can be observed,
adsorption of Hydrogen and hydroxyl species can be found on a SiO2 substrate
and be observed using HREELS (see Fig 5.12).
SiO2/Si
Intensity (a.u.)
H
O
O
Si
Si
O
Si Si
H
O
O
Si
H
H
O
C
Si
0,05 0,10 0,15 0,20 0,25 0,30 0,35 0,40 0,45 0,50
Electron energy loss (eV)
Fig 5.12 HREELS spectrum of a SiO2 surface before annealing .Vibrational bonds of O, H, OHabsorbed on the surface are observed. By D. Stolcic et. al.,[120].
In order to image Au nanoparticles on Silica layers, STM measurements
were taken. It turned out that the tunnelling voltage between tip and surface,
necessary for obtaining STM images, is significantly higher than in the case of
HOPG. This is because of the semiconducting properties of SiO2 films.
Furthermore surface roughness of SiO2 thin films is higher than those of HOPG
because of the underlying polycrystalline Si substrates. Obtaining accurate
STM images is difficult, since the higher the quality of the image is taken, the
less the roughness of the surface should be. Eventually, it is worth to mention
44
5 Experimental Setup
that the interaction of particles with the SiO2 film is stronger in comparison with
the HOPG substrate, avoiding particles sintering on the SiO2 film surface.
6 Results and Discussion
6.1 Ag and Au Nanoparticles on HOPG
Ag growth modes on mild- and heavily sputtered HOPG were studied by
means of STM and subsequently XPS. These results were compared with those
of Au nanoparticles on the same substrate. Both systems were prepared in the
same way, and therefore, the results could be directly compared. Moreover, Ag
(or Au) nanoparticles on HOPG were compared with growth behaviors of these
particles on WSe2 [121]. Although both, HOPG and WSe2, are Van-der-Waals
surfaces, different growth modes were found. This indicates that subtle changes
in metal-support interaction can alter particle shapes considerably.
STM Analysis
Ag Nanoparticles on Sputtered HOPG
In Figures 6.1-6.4, Constant Current Topography (CCT) STM images of
four different Ag/HOPG surfaces are presented. These images were taken under
UHV conditions at room temperature. The particle diameter was determined by
measuring the Full Width at Half-Maximum (FWHM) of the particle profile.
The defect density was controlled by altering the sputtering time, in order to
study the influence of defects in the substrate. Different particle densities and
sizes were obtained not only by varying the density of defect sites but also by
controlling the evaporation time. The emission current24 was maintained
constant at a value of 15 mA. The coverage was estimated with a combined
STM/XPS quantitative analysis, which is extensively explained later in this
chapter (see section of this chapter: “Combined STM/XPS quantitative analysis
of the lateral particle size”).
Fig 6.1 shows a STM image of Ag nanoparticles prepared by sputtering
HOPG for 5 seconds at an Ar pressure of 5 x 10-5 mbar and evaporating Ag for
8 minutes at room temperature. The average particle size amounts to 3 nm in
diameter and 1 nm in height. Particles show oblate forms with larger diameters
24
The emission current is related to electrons emitted from the filament to the metal target.
45
46
6 Results and Discussion
than heights, indicating that particles grow more laterally than normal to the
surface in this small size regime. This kind of growth was already found for
metal particles on other substrates [15, 72]. A convolution effect of the tip
cannot be completely excluded as a part of the explanation for this oblate form.
•Defect density: ~ 2.5 % of a ML
•Deposition rate: ~ 0.6 ML/min
•Coverage: ~7 %, ~ 5 ML
•Particle density: ~104 µm-2
60
50
Count
Diameter: ~ 3 nm
30
20
Count
50
40
40
Height: ~ 1nm
30
20
10
10
0
1
2
3
4
5 6
7 8
9 10 11 12 13 14 15 16
Particle diameter (nm)
0
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle height (nm)
Fig 6.1 Two STM images of the same sample for Ag nanoparticles on HOPG; Upper left:
135.3 x 132.3 nm2. Upper right: 50 x 50 nm2. STM parameters: 3.0 V, 0.7 nA. Lower pictures:
Diameter and height distribution of the particles.
In Fig 6.2, Ag particle size as well as particle density increases after
sputtering 5 seconds at an Ar pressure of 5 x 10-5 mbar and evaporating Ag for
30 minutes. The average diameter and height of the particles is about 7 nm and
3 nm, respectively.
6 Results and Discussion
•Defect density: ~2.5 % of a ML
•Deposition rate: ~ 0.5 ML/min
47
•Coverage: ~10 %, 15 ML
•Particle density: ~ 104 µm-2
160
80
140
120
60
Count
Count
Diameter: ~ 7 nm
40
Height: ~ 3nm
100
80
60
20
40
20
0
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle diameter (nm)
0
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle height (nm)
Fig 6.2 Two STM images of the same sample for Ag nanoparticles on HOPG; upper Left: 600x
320 nm2. Upper right: 214.5 x 214.5 nm2. Tunneling parameters: 1.5 V, 1.0 nA. Lower figures:
Diameter and height distribution of the particles.
In Fig 6.3, the sputtering time was reduced to 2 seconds under the same
Ar conditions and afterwards, Ag was evaporated for 60 minutes. Mean particle
size accounts to 6 nm in diameter and 2 nm in height. Smaller particle sizes can
be observed, even though the Ag exposure was increased by a factor of 2
compared to Fig 6.2. This indicates that the reduced sputtering time decreases
the sticking probability. Hence, an influence of sputtering time on particle
growth can be directly observed.
48
6 Results and Discussion
•Defect density: ~1 % of a ML
•Deposition rate: ~ 0.2 ML/min
•Coverage: ~18 %, ~ 10 ML
•Particle density: ~ 104 µm-2
60
50
80
70
Diameter: ~ 6nm
60
40
50
Count
count
Height: ~ 2 nm
30
20
40
30
20
10
0
10
0
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle diameter (nm)
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle Height (nm)
Fig 6.3 Two STM images of the same sample of Ag nanoparticles on sputtered HOPG; upper left: 352.3
x 352.3 nm2. Upper right: 154.9 x 154.9 nm2. STM parameters: 2.0 V, 2.4 nA. Lower pictures: Diameter
and height distribution of the particles.
Another HOPG sample was sputtered for 2 seconds and Ag was
evaporated for 90 minutes (See Fig 6.4.). Average diameter and height of
particles are 10 nm and 5 nm, respectively. The particles show not perfectly
circular shape and faceted nanoparticles are detected. The particle shape is also
rather similar to those of Au nanoparticles on HOPG found by our group [117].
6 Results and Discussion
49
•Defect density: ~1 % of a ML
•Deposition rate: ~ 0.3 ML/min
•Coverage: ~15 %, 25 ML
•Particle density: ~ 104 µm-2
60
Diameter: ~ 10 nm
100
Count
Count
50
120
40
30
80
60
20
40
10
20
0
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle diameter (nm)
Height: ~ 5nm
0
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle Height (nm)
Two STM images of Ag nanoparticles on sputtered HOPG; Upper left: 394.4 x 349.8
nm2. Upper right: 45 x 45 nm2. STM parameters: 1.9 V, 2.1 nA. Lower picture: Diameter and
height distribution of the particles.
Fig 6.4
A three-dimensional growth of the particles is observed, when one
compare the growth of the 4 sample of Fig 6.1-Fig 6.4. This is different from
the growth behavior observed for Au and Ag nanoparticles on WSe2 by
Leiderer et. al. [121]. In this work, a perfect two-dimensional growth of the
particles is preferred, leading to the formation of particles of 2 nm in height and
larger than 20 nm in diameter. A similar two-dimensional growth was also
found for Au clusters grown on nanometer–sized pits created on a graphite
surface at higher temperatures (600 K) [122].
This three-dimensional growth is shown in detail in Fig.6.5, in which
particle diameter versus particle height is plotted. In Fig 6.5, one can observe
that smallest particles grow more laterally than normal to the surface, while
particle growth perpendicular to the surface is commonly observed for larger
particles.
50
6 Results and Discussion
height (nm)
A summary of the parameters used for the preparation of the Ag
nanoparticles grown on HOPG is presented in Table 1.
12
11
10
9
8
7
6
5
4
3
2
1
0
Ag growth Mode
0
1
2
3
4
5
6
7
8
9 10 11 12
diameter (nm)
Fig 6.5
Particle diameter versus particle height of Ag nanoparticles grown on HOPG.
Ag
Diameter
Height
Coverage
(ML, %)
Deposition
Rate
(ML/min)
Particle
Density
Defect
Density
(% of ML)
Fig 6.1
3 nm
1nm
5 ML, ~7 %
0.6min-1
~104 µm-2
2.5%
Fig 6.2
7 nm
3nm
15ML,~10%
0.5min-1
~ 104 µm-2
2.5 %
Fig 6.3
6 nm
2 nm
10ML,~18%
0.2 min-1
~104 µm-2
1%
Fig 6.4
10 nm
5 nm
25ML, 15%
0.3 min-1
~104 µm-2
1%
Table 1 Estimation of the parameters used for the preparation of Ag nanoparticles grown on
HOPG. Note: see section “Combined STM/XPS quantitative analysis of the lateral particle
size”.
6 Results and Discussion
51
Au Nanoparticles on Sputtered HOPG
In this section the growth of Au nanoparticles on HOPG with different
densities of defect sites and different size regimes is presented. Au particles
were prepared in an analogous manner as Ag particles: First, creating defect
sites by sputtering, which serve as nucleation centers for metal atoms, and
afterwards, evaporating Au atoms onto sputtered HOPG substrate.
The sample of Fig 6.6 was prepared by sputtering the HOPG surface for
10 seconds at an Ar pressure of 5 x 10-5 mbar. Afterwards, Au was evaporated
for 2 minutes with an emission current of 15 mA. Particles have a diameter of
about 5 nm and a height approx. of 23 nm.
•Defect density: ~5 % of a ML
•Deposition rate: ~ 5 ML/min
•Coverage: ~13 %, 10 ML
•Particle density: ~ 104 µm-2
25
50
20
40
15
30
Counts
Count
Diameter: ~ 5 nm
10
5
0
Heitght: ~ 2-3 nm
20
10
0 1 2 3 4 5 6 7 8 9 10111213141516
Particle diameter (nm)
0
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle Height (nm)
Fig 6.6 STM image of a sample for Au nanoparticles on HOPG :( 300 x 300 nm2). Tunnelling
parameters: -5.3 V, 0.08 nA. Lower picture: diameter and height distribution of the Au
particles.
52
6 Results and Discussion
Fig 6.7 shows 2- and 3-D STM images of a sample, which was sputtered
for 5 seconds under the same Ar conditions. Afterwards, Au was evaporated at
an emission current of 17 mA during 5 minutes. By decreasing sputtering time
and increasing the amount of Au atoms deposited on the surface, the density of
defect sites decreases and larger Au particles appear. Closer examination of the
Au islands reveals no significant change in height, which may imply that 2
dimensional growths dominate in this smaller size regime. The lateral size of
the particles is about 8 nm and the height is about 3 nm. However, the observed
height can be subject to a relative error, since an extremely high sample bias of
about 2 V was taken to display these images. The high voltage was needed in
order to obtain a stable image.
•Defect density: ~2.5 % of a ML
•Deposition rate: ~ 3 ML/min
70
Diameter: ~ 8 nm
25
60
20
50
counts
Count
30
•Coverage: ~18 %, ~ 12 ML
•Particle density: ~ 104 µm-2
15
10
40
30
20
5
0
Height: ~ 3 nm
10
0 1 2 3 4 5 6 7 8 9 10111213141516
Particle diameter (nm)
0
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Particle Height (nm)
Fig 6.7 Two STM images of the same sample for Au nanoparticles on HOPG; Upper left: 3-D
image (256.5 x 256.5 nm2). Upper right: 2-D image (256.5 x 256.5 nm2). Tunnelling
parameters: -1.97 V, 0.08 nA. Lower pictures: Diameter and height distribution of the particles.
6 Results and Discussion
53
When the density of defects is decreased by means of sputtering for only
3 seconds at an Ar pressure of 5 x 10-6 mbar and Au is evaporated for 15
minutes with a higher current emission of 22.5 mA, lateral as well as vertical
dimensions of particles increase drastically (see Fig 6.8).
•Defect density: ~1.5 % of a ML
•Deposition rate: ~ 3.3 ML/min
20
•Coverage: ~24 %, ~ 50 ML
•Particle density: ~ 102 µm-2
Diameter: ~ 40 nm
30
Count
count
Height: ~ 10 nm
25
15
10
5
20
15
10
5
0
20
30
40
50
Particle diameter (nm)
60
0
4 5 6 7 8 9 10 11 12 13 14 15 16
Particle height (nm)
Fig 6.8 Two STM image of the same sample for Au nanoparticles on HOPG; Upper left: 430 x
430 nm2; tunneling parameters: 4.0 V, 0.22 nA. Upper right: 95 x 145 nm2; tunneling
parameter 2.2 V, 0.22 nA. Lower picture: Diameter and height distribution of the particles.
Note: Existence of steps between different terraces (111) is observed.
A three dimensional growth mode with Au particles of about 40 nm in
diameter and about 10 nm in height is observed. The growth mode of these
particles is definitely in line with the growth mode previously observed for Ag
nanoparticles on HOPG, but in contrast to the 2 dimensional growth mode
54
6 Results and Discussion
reported for Ag and Au nanoparticles on WSe2 at this particle size regime. Fig
6.8 reveals facet structures in single particles, which are similar to those
observed in Ag nanoparticles grown on HOPG. Further images show that these
structures can be observed more clearly at higher Au coverages.
This is the case of the sample in Fig 6.9. This figure shows the formation
of nanostructured Au thin film instead of separate particles. In this case, the
HOPG surface was sputtered for 5 seconds at an Ar pressure of 5 x 10-6 mbar
and subsequently, Au was evaporated for 10 minutes, using an emission current
of 22.5 mA. Another STM image taken from the same sample is displayed in
Fig 6.10. The existence of steps on nanoparticles can be identified. The
particles show hexagonal planes with preferred formation of the (111) plane
parallel to the surface.
STM image of Au nanoparticles (nanostructured film) grown on HOPG (350 x 500
nm2). Tunneling parameters: -1.2 V, 1.3 nA.
Fig 6.9
6 Results and Discussion
55
13.8 nm
17.4 nm
1,8
1,8
1,6
nm
nm
1,6
1,4
1,4
1,2
1,2
1,0
0
2
4
6
8
nm
10
12
14
0
2
4
6
8
10
12
14
16
18
nm
Fig 6.10 Upper picture:STM image of Au nanoparticles grown on HOPG (220 x 220 nm2). Tunnelling
parameters: -1.2 V, 1.3 nA. Lower picture: Step profiles of about 0.4 nm in height between different
terraces are observed.
Fig 6.11 displays the growth mode of Au on HOPG, in which particles
undergo a three dimensional growth with increasing particle size. A summary
of the parameters used for the preparation of the Au nanoparticles for the Figure
6.6-6.8 is presented in Table 2.
56
6 Results and Discussion
45
Height (nm)
40
35
Au Growth Mode
30
25
20
15
10
5
0
0
5
10
15
20
25
30
35
40
45
Diameter (nm)
Fig 6.11 Au particle diameter versus particles height.
Au
Diameter
Height
Coverage
(ML, %)
Deposition
Rate
(ML/min)
Particle
Density
Defect
Density
(% of ML)
Fig 6.6
~5 nm
~2 nm
10 ML, 13 %
5 min-1
104 µm-2
5%
Fig 6.7
~8 nm
~2-3 nm
12 ML, 18 %
3 min-1
104 µm-2
2.5 %
Fig 6.8
~40 nm
~10 nm
50 ML, 24 %
3.3 min-1
102 µm-2
1.5 %
Table 2 Estimation of the parameters used for the preparation of Au nanoparticles grown on
HOPG. Note: see section “Combined STM/XPS quantitative analysis of the lateral particle
size”.
6 Results and Discussion
57
Combined STM/XPS quantitative analysis of the lateral particle size
Determination of the particle height using STM may contain errors of
several tens of nm, since the “apparent particle height” is a result of electronic
as well as geometric contrast [72, 113]. Generally, the error in lateral particle
size is more important, yet an overestimation due to the limited lateral
resolution of STM tips, particle sintering, etc is always involved. This
overestimation can be corrected by combining STM measurements with the
quantitative analysis of XPS data. From this analysis one can determine the
correction factor for the “apparent lateral particle size by means of STM, which
is generally larger than the real lateral size. It is well known from earlier studies
by Hövel et al. [113], that the Full Width at Maximum Height (FWMH) of a
particle can overestimate the particle size by ~60 % compared to TEM
estimations. Taking this information into account, some calculations were
carried out. In order to calculate Ag and Au coverages on HOPG for different
samples, the percentage of the HOPG surface, covered by metal nanoparticles,
was estimated by means of STM. The number of particles in a unit area was
determined and then multiplied by the FWHM, assuming spherical particles
(see Fig 6.12).
1,2
1,0
nm
0,8
2.2 nm
0,6
0,4
0,2
0,0
0
2
4
6
8
10
12
14
nm
Fig 6.12 Left: STM image of Ag nanoparticles grown on HOPG (50 x 50 nm2) Tunneling
parameters: 3.0 V, 0.7 nA. Right: FWHM of the particle profile is shown in the plot. In this
case, the STM estimation of the particle size is about 2.2 nm. Note that a step of about 0.6 nm is
located at the left lower corner of the STM image.
The coverage was also calculated by means of an XPS analysis. In case of
Ag, the intensity ratio I[Ag(3d5/2)]/I[C(1s)] for the corresponding peaks was
58
6 Results and Discussion
determined. Note that the cross section of the C 1s state in HOPG per unit
volume is about two times larger than that of Ag 3d5/2; in other words; a pure
Ag sample generally shows a higher intensity of Ag 3d5/2 states than the C 1s
intensity of a pure HOPG by a factor of 2:
I0C 1
=
I0Ag 2
(1)
where I0Ag and IC0 are the intensity for pure Ag and C surfaces,
respectively. It is worth to mention that the experimental setup reproduces the
absolute intensities of the photoemission peaks within an error bar of about
10% [100].
To explain this analysis, an important idea about XPS must be illustrated.
XPS is a surface sensitive technique, i.e., surface topmost layers contribute
more to the photoemission signal than deeper layers [100]. With increasing Ag
film thickness, one can assume an exponential decrease and an increase of the C
and Ag peak intensities, respectively. Moreover, assuming that 2 nm is the
mean free path λ of the photoelectrons from C 1s and Ag 3d [100], the intensity
of the C 1s state, IC, may be composed of contributions from the bare HOPG
surface and an attenuated contribution due to the part of the surface covered by
Ag:
⎛ −d
⎞
I C = I C0 (1 − ϑAg ) + ϑAg I C0 exp ⎜ Ag ⎟
λ
⎝
⎠
(1)
where HOPG is covered to a fraction ϑAg with deposited Ag
nanoparticles, dAg is the thickness of the particle previously estimated by STM
and λ is the mean free path of the photoelectrons.
Therefore the intensity of the C 1s state IC is given by
⎡
⎛ -d
IC = I0C ⎢1 − ϑAg + ϑAg ⋅ exp ⎜ Ag
λ
⎝
⎣
⎞⎤
⎟⎥
⎠⎦
(2)
and the signal from the deposited Ag particles on HOPG is given by
⎡
⎛ d ⎞⎤
I Ag = ϑAg ⋅ I0Ag ⎢1 − exp ⎜ − Ag ⎟ ⎥
⎝ λ ⎠⎦
⎣
(3)
6 Results and Discussion
59
Consequently,
⎡
⎛ d Ag ⎞ ⎤
⎟⎥
I Ag I0C
⎝ λ ⎠⎦
⎣
⋅
=
IC I0Ag
⎡
⎛ d
1 − ϑAg + ϑAg ⎢exp ⎜ − Ag
⎝ λ
⎣
ϑAg ⎢1 − exp ⎜ −
⎞⎤
⎟⎥
⎠⎦
(4)
The fraction ϑAg , in which HOPG is covered with Ag nanoparticles, is then
given by
ϑAg
⎛ I Ag 1 ⎞
⋅ ⎟
⎜⎜
IC 2 ⎟⎠
⎝
=
⎡ ⎛ I Ag 1 ⎞ ⎤ ⎡
⎛ d ⎞⎤
⋅ ⎟ ⎥ ⋅ ⎢1 − exp ⎜ − Ag ⎟ ⎥
⎢1 + ⎜⎜
⎟
⎝ λ ⎠⎦
⎢⎣ ⎝ IC 2 ⎠ ⎥⎦ ⎣
(5)
For the specific case of Ag particles from the sample of Fig 6.4, Ag
nanoparticles have a mean diameter of 10 nm and a mean thickness of about
5nm. From our STM image in Fig 6.4 about 40-45% of the HOPG surface is
estimated to be covered by Ag. The intensity ratio I[Ag (3d5/2)]/I[C (1s)]
obtained by means of XPS is 0.27 for this sample. Taking into account equation
(5), the estimation of HOPG covered with Ag nanoparticles is about 13 percent.
This value, as real coverage ratio, corresponds to the situation in which about
15% of the HOPG surface is covered by a 5 nm thick Ag film. If we compare
with the STM estimation of ~ 45% of Ag coverage, the HOPG area covered by
the particles is overestimated by about a factor of 3. Therefore, the real particle
diameter should be about 60 % of the particles size estimated by STM.
Analogous results can also be found for the sample in Fig 6.2. In the STM
image, Ag nanoparticles have an apparent particle diameter of 7 nm and about
30 % of the surface seems to be covered by Ag (with a mean particle thickness
of 3 nm); however, the intensity ratio I[Ag(3d)]/I[C(1s)] amounts to 0.07,
implying that only 10% of the HOPG surface is covered by a 3 nm thick Ag
film. Again the real particle diameter is only 60 % of the apparent particle size
in the STM image. For the sample of Fig 6.1, an overestimation by a factor 3
turns out after calculations once again. With Au nanoparticles, same
considerations can be taken into account. Note that the photoemission cross
sections of the Ag 3d and Au 4f states differ only less than 10% [100]. The
Au/HOPG sample of the STM image in Fig 6.7 yields an intensity ratio
60
6 Results and Discussion
I[Au(4f)]/I[C(1s)] comparable to that of Ag(3d)/C(1s) for the Ag/HOPG sample
with a equivalent particle size and density, so that Fig 6.7 can be considered to
be a representative image of the Au/HOPG sample. In the next part of this
chapter (XPS Studies), particles sizes referred in the XPS Spectra have been
previously corrected by means of this combined XPS/STM estimation.
Particle Sintering
In the present work, particle sintering was sometimes observed during
STM measurements. This phenomenon can affect coverage estimation by STM
and XPS in a different way, since the STM tip scans in a very small area (~1µm
x 1µm), whereas XPS affects a larger area (~ 1mm x 1mm).
Two examples of particle sintering for a non-sputtered HOPG substrate
are given in Fig 6.13. Wide vacant areas with large Ag islands of approximately
8-10 nm diameters and about 2 nm height can be observed on the left image in
Fig 6.13. In contrast, particles inside the marked square area in blue undergo no
sintering. Here, particles are about 4 nm in diameter and 2 nm in height. The
right image in Fig 6.13 shows another STM image of the same sample taken
some minutes later at another location. High concentration of particles with
desert areas were displayed, which indicate that tip interaction, during scanning,
yields this sintering.
Fig 6.13 Two STM images of the same sample for Ag nanoparticles on non-sputtered HOPG. The images
were taken at different spots of the sample. Left: 500 x 500 nm2. Right: 495 x 335 nm2. Tunnelling
parameters: -1.07 V, 0.06 nA. Note: Particles inside the marked square area undergo no sintering.
6 Results and Discussion
61
Fig 6.14 displays three time-successive STM images, in which the
mobility of particles deposited via evaporation of Ag metal atoms on nonsputtered HOPG is shown. An inhomogeneous particle size distribution can be
observed. This suggests that islands are highly mobile on the surface at room
temperature and undergo agglomeration. Particles formed not only by
evaporation but also by particle diffusion on the surface appear more frequently
at step edges than somewhere else. Weak metal-support interaction between
substrate HOPG and Ag metal nanoparticles explain this fact.
a
b
c
Fig 6.14 Time successive STM images at same location for a HOPG sample on which Ag
nanoparticles were grown. Tunnelling parameters: -3 V, 0,7 nA. a): 150 x 115 nm2; b): 146 x
114 nm2 (6 minutes after a); c): 147 x 109 nm2 (8 minutes after b).
Particle sintering usually happens, when a STM tip draws near enough to
the surface to measure a sufficient tunneling current. Tunneling parameters
should be controlled and varied more or less randomly in order to stop further
approaching of the tip to the surface, creating some disturbance of the particle-
62
6 Results and Discussion
tip interaction. Another reason might be that STM tips are contaminated by
metal particles, which perturb the imaging process. Eventually, it can also occur
that changes of the tunneling parameter at initial stage of STM measurement
can cause movement and agglomeration of Ag particles. It should be mentioned
that scan speed (300 nm/s in case of sample in Fig 6.14) can also affect the
sintering process.
Ag particle sintering on HOPG surfaces can be excluded in case for the
samples in Fig 6.1, Fig 6.2 and Fig 6.4, since the coverage ratio Ag/C estimated
by STM is higher than the one of the XPS data. This comparison between STM
and XPS data does not correspond with agglomeration, because sintering would
have made the surface/volume ratio decreasing in all of these cases. It is
important to note that the images do not change during several hours of
scanning. These features usually appear predominantly at step edges by nonsputtered HOPG, indicating that defect sites created by sputtering HOPG may
play an equivalent role for the mobility of the nanoparticles.
Comparison of Ag and Au nanoparticles on sputtered and non-sputtered
HOPG
An increase of the sticking probability of Ag or Au atoms by creating
defect sites on HOPG is obvious, since STM results show that average particle
size becomes larger with an increased density of the defect sites, created by
applying the same evaporation time.
Two STM images of two different samples are displayed in Fig 6.15.
Those two samples were prepared with the same Ag evaporation time on
sputtered and non-sputtered HOPG surfaces. The images in Fig 6.15 show that
particle density and particle size are reduced when the surface is not sputtered.
Two STM images of the same sample, in which Ag atoms were evaporated on
non-sputtered HOPG, are shown in Fig 6.16. Note that nanoparticles
preferentially form at step edges and kinks. Low particle density can be
observed in Fig 6.16 in comparison to Ag particle density on sputtered HOPG,
in which a decoration of particles on step edges is more difficult to find.
The sticking probability of metal atoms on a non-sputtered HOPG was
estimated to be about 0.1. To obtain this value, the sticking probability in a
sputtered HOPG surface was normalized to one as reference. For this
6 Results and Discussion
63
estimation, Ag coverage ratios were compared, using STM images of the
samples displayed in Fig 6.15.
Fig 6.15 Left: STM image of Ag nanoparticles on sputtered HOPG (214.5 x 214.5 nm2); Tunneling
parameters: -1.5 V, 1.0 nA. Right: Ag nanoparticles on non-sputtered HOPG with same Ag exposure time
(30 min); (207.2 x 181.4 nm2); Tunneling parameters: -2V, 2.0 nA.
Fig 6.16 Two STM images of the same sample of Ag nanoparticles on a non-sputtered HOPG
surface, demonstrating that Ag preferentially decorates step edges. Left: 871.1 x 871.1 nm2,
Right: 246 x 246 nm2. Tunneling parameters: 0.2 V, 1 nA.
STM imaging of metal nanoparticles on non-sputtered HOPG can be
affected by noise due to sintering, tip contaminations and nucleation of bigger
clusters. Considering the significant increase of the sticking probability of the
metal atoms on HOPG surfaces after sputtering, one can argue that the surface
is rather defective by means of sputering. However, the HOPG surface plasmon
64
6 Results and Discussion
resonance of C1s state could still be identified by our XPS spectra, on which a
broadening of the C 1s peak was observed, indicating that only some defect
sites are created in the surface by sputtering without high changes of the
roughness.
Morphology and Growth of Au and Ag Nanoparticles on HOPG
In this section, the morphology and growth of Ag and Au metal
nanoparticles on HOPG is described by means of analyzing STM images. STM
technique in situ can be very useful for these kind of studies, since submonolayer sensitivity over large fields of view (~1 µm2) and the possibility to
image large variations of cluster density depending on substrate, deposition
technique, temperature etc, can provide detailed information about the kinetic
properties of the system. As a result of the geometry and the growth mode of a
particle on a substrate, information about the kinetic processes involved in the
formation of a particle can be obtained.
Three growth modes for a material deposited on a surface have been
suggested a long time ago by Bauer 1958 [123]. The Frank-Van Der Merwe
mode (layer by layer growth) arises because the atoms of the deposited material
are more strongly attracted by the substrate than they are to themselves. When
the opposite case takes place, Volmer-Weber mode (island growth), atoms
seem to be more strongly bound to each other than to the substrate, forming
islands. The intermediate case, Stranski-Krastanov mode (Layer-plus-island
growth) happens when layers are formed first, but then the system changes its
growth behavior and continues in the island-growth mode (see Fig 6.17).
In our case, the island growth mode dominates the growth of Ag and Au
nanoparticles on HOPG. An example of island growth is shown in Fig 6.18 in
which a three dimensional STM image of Ag nanoparticles on HOPG is
displayed.
The formation of nanoparticles can be understood via diverse atomistic
surface processes [80]. Fig 6.19 shows some of these processes like adsorption
and desorption of metal atoms at the surface, terrace diffusion, island formation
via nucleation within each layer; rapid diffusion at island edges of atoms, step
formation or interlayer diffusion through step-edge barriers, etc.
6 Results and Discussion
Layer
growth
65
Layer plus island growth
Island growth
θ<1ML
1<θ<2ML
θ<2ML
Fig 6.17 Schematic diagram of the three growth modes usually observed for metal overlayers
[80].
Fig 6.18 Three dimensional STM image of Ag nanoparticles on HOPG with a particle size distribution of
3 nm in diameter and 1 nm in height. (50 x 50 nm2). Tunneling parameters: 3.0 V, 0.7 nA.
Absorption
Desorption
Step formation
Special sites
Surface
Nucleation
Interdiffusion
Fig 6.19 Atomic surfaces processes and characteristic energies, which participate in nucleation
and film growth.
66
6 Results and Discussion
From an energetic point of view, the growth mode of a metal on a surface
is a consequence of minimizing the surface free energies25 of metal, substrate
and interface [124, 125]. The surface free energy of noble metals (e.g.; Ag ~1.3
J/m2) [126, 127] is more than 3 times higher than the surface free energy for a
Van der Waals substrate like HOPG (~ 0.3 J/m2). Therefore, as long as the
surface does not undergo a sputtering process, both atoms, Ag and Au
respectively, can diffuse easily on the surface overcoming the barriers that
separate the binding sites created by no defective benzene rings. The left picture
in Fig 6.20 shows a representation of the diffusion barrier in one dimension as a
function of the position for an ideal defect-free surface. Note that the existence
of an analogy between periodicity of the attractive potential wells (left, Fig
6.20) and spatial periodicity of the STM image with atomic resolution of a
defect-free HOPG surface (right, Fig 6.20), is evident. As already mentioned,
Au and Ag nanoparticles on HOPG undergo a 3 dimensional growth with
increasing particle size. Particles grow, exhibiting polyhedral shapes, mostly
hexagonal with facet formation (see Fig 6.21).
Fig 6.20 Left: One dimensional representation of the diffusion barrier as a function of position
on an ideal defect-free surface. Right: Geometrical corrugation shown by a three dimensional
STM image with atomic resolution of a defect-free HOPG surface (1.7 x 1.0 nm2; 0.1 V, 0.7
nA).
These facet structures on nanoparticles were extensively explained with
energetic concepts, which suggest that a “three dimensional equilibrium crystal
shape is governed by the orientation-dependent of a surface energy γ [128]. The
minimum value in γ corresponds to a small number of different facets.
Quantitative information on surface energies in facets of supported clusters
were reported in previous results by F. Besenbacher , H.-J. Freund et al., [129].
25
The surface free energy is defined as the work required for increasing the area of a substance
by one unit area.
6 Results and Discussion
67
The existence of many steps with atomic height between terraces (~0.2
nm) was found for Au nanoparticles on HOPG as shown in Fig 6.22.
b
(111)
(100)
(111)
c
a
2
Fig 6.21 a) STM image of 2 Au nanostructured particles (60 x 60 nm ). Tunnelling parameters:
1.5V, 0.2 nA. b) Schematic draw of the facet structure of the particle; Hexagonal terraces with
plane (111) connected with step planes (100) and (111) form the particles. c) Schematic draw of
atoms diffusing on the surface subject to an Ehrlich-Schwöbel barrier. The potential energy
schema of the surface is shown below.
Particle Profile
16.4 nm
2,0
nm
1,5
0.2 nm
1,0
0,5
Atomic layer step: ~ 0.2 nm
0,0
0
2
4
6
8
10
12
14
16
nm
Fig 6.22 Left: STM image of a single Au nanoparticle (44 x 44 nm2). Tunnelling parameters:1.5
V, 0.2 nA. Right: Step profile of the blue line. Atomic steps of ~0.2 nm size can be recognized.
68
6 Results and Discussion
The formation on the particle of terraces (111) perpendicular to the
substrate (HOPG) can be explained, taking into account the Ehrlich-Schwöbel
barrier26 (see Fig 6.21 c). Metal atoms landed on the topmost terrace of a
particle cannot diffuse down to the next atomic layer, due to this barrier,
forming an additional plane on the top of the terrace (111). Some methods
developed for the evaluation of this step edge barrier related to interlayer
transport in metal island growth were reported by Markov et al [130, 131].
In contrast to our work, Hövel et al. deposited Au nanoparticles on HOPG
at elevated temperatures of 600 K, creating nanopits on a HOPG surface via
sputtering, and the formation of an ideal truncate octahedral particle shape was
found [113]. In our experiments, non-ideal truncated octahedral particles were
observed. Kinetic arguments can justify the difference between our results and
those from Hövel and coworkers. The diffusion barrier across step-edges has
been determined to be 0.06-0.07 eV with a prefactor of 1012 s-1, implying that
the shapes of particles grown at room temperature and higher temperatures
(e.g., 600 K) can be much different. A deposition at room temperature is not
sufficient to overcome the activation barrier (Ehrlich-Schwöbel barrier) of the
metal atoms jumping over the step edges.
Beyond kinetic factors, the same ideal truncated octahedral shapes of Au
and Ag nanoparticles on the Van der Waals surface WSe2 at room temperature
was observed by P. Leiderer et al., [121]. In this work, the step structure has
never been identified. Ag and Au on WSe2 initially grow in a three dimensional
mode; nevertheless, as the particle height becomes about 2 nm, almost a two
dimensional growth of the particles is preferred, in which flat metal particles
with a lateral size of 20 nm and a height of 2 nm were observed. This indicates
that, even for Au films thicker than 10 atomic layers (2-3 nm), growth of metal
nanoparticles (or films) can be influenced by the underlying substrate. Both
WSe2 and HOPG are Van-der-Waals surfaces, however, subtle differences in
the metal-support interactions between these substrates seem to have influence
on the particle growth.
Against kinetic and metal-support considerations, one may argue that the
deposition rate of atoms in our experiment was too fast, producing these kind of
stepped structures at room temperature. When the deposition rate is too fast, Au
and Ag atoms could form clusters on the topmost terrace before moving down
to the next underlying plane, because clusters cannot diffuse as easily down to
26
Step edge energy barrier for the interlayer diffusion of adatoms.
6 Results and Discussion
69
the atomic layer as metal atoms. To check whether the deposition rate does
actually influence the shape of Au nanoparticles, the same amount of Au as in
Fig 6.22 with a deposition rate reduced by a factor of about 20-30, was
evaporated on sputtered HOPG. The results were the same as in our previous
experiments. The step structure could not be found, hence, the hypothesis that
these results are due to the deposition rate can be excluded.
The particle composed of facets, shown in Fig 6.22, looks crystalline in
shape. Furthermore, a loss of resolution is clearly observed at the kinks. In
general, the sensitivity to determine particle shapes using STM is limited, since
STM imaging usually produces some ‘‘artificial’’ smoothing of features. This
smoothing effect was predicted via simulation of STM images by K. J.
Caspersen et al. [132]. A qualitative explanation for this fact is illustrated in Fig
6.23. At the kinks of the top facets (111) and lateral facets (100), the tunneling
current changes, since the outermost apex atom from the tip cannot scan
anymore beyond the end of the particle. Instead of this apex, some others atoms
somewhere on the side of the tip, which has not achieved the kink yet, begin to
scan, resulting in a loss of resolution. For this reason, one often assumes that
particles have ideal forms such as truncated octahedron shapes, even though it
may be only due to the STM resolution. Our results suggest that particles do not
form truncated structures and they can have a much higher defect density (e.g.
layer steps) than estimated under the assumption of an ideal particle form.
Outer Apex of a STM tip
Scan direction
kink
(100)
(111)
Substrate
Fig 6.23 Illustration of a kink of a nanoparticle formed by facets (100) and (111). Above these
particles a STM tip with two apexs scans the particle from right to the left.
70
6 Results and Discussion
It is well-known that step edges can often be active sites for
heterogeneous catalysis, pointing out that the chemical activity of these
nanoparticles can be related to the particle geometry [63]. Therefore, a good
description of the morphology of these particles is important. Our results show
that there is no significant difference in the particle morphology and growth
between Ag and Au on sputtered HOPG, where facet formation, hexagonal
shapes of particles and a growth with an axial ratio in direction (111) were
observed.
Thitherto only geometric arguments were presented. There are two important
questions to be considered: First; how the geometry of particles is related to
the electronic structure of those nanoparticles, and second; to what extend are
geometry and electronic structure affected by metal-support interaction. To
gain a better insight into these issues, the electronic structure of these
particles was studied using XPS.
6 Results and Discussion
71
XPS Studies
Core level binding energy shifts with decreasing particle sizes of Au and Ag
nanoparticles grown on HOPG were studied, using X-ray Photoelectron
Spectroscopy, (XPS). The results of Ag metal particles were compared to
those of Au nanoparticles. In terms of final and initial state effects, the origin
for the observed size-dependent binding energy shifts is discussed
extensively in this section.
Ag 3d Core Level Spectra as a Function of Particle Size
Fig 6.24 a) shows 4 different XPS spectra of the Ag 3d core levels for the
samples in Fig 6.1-6.4 already imaged by STM. Each peak was fitted by using a
combination of 2 Gaussian functions. A main peak and a shoulder at higher
binding energies have been combined to determine the binding energies of the
Ag 3d states accurately. Positive core level shifts of Ag 3d states with
decreasing particle size with respect to the Ag bulk value (368.3 eV) are
observed. For the smallest particle with a diameter of about 2 nm, a core level
shift of +0.6 eV is accounted with respect to the 3d levels of the Ag bulk
crystal. Comparing with bulk core levels, typical chemical shifts between 0.6
and 0.8 eV have been previously observed for diverse metal nanoparticles [133140].
In Fig 6.24 b), binding energies and widths (FWHM) of the Ag 3d states as a
function of the particle size are summarized. In combination with STM
images, the corresponding energy shifts of the Ag 3d spectra relative to the
bulk value enable us to estimate the particle diameter. It is worth mentioning,
that FWHM increases with a decreasing of particle size. The particle size
distribution observed in STM images becomes broader with lower particle
size. Hence, this indicates that broadening of Ag 3d core levels may be
correlated to the particle size distribution. In addition, It should be related to
different life-time broadenings of the hole state by the screening process27 or
due to a possible charge transfer between nanoparticles and substrate. These
mechanisms are of statistical nature.
27
Screening process: Final state relaxation of the positive hole in a cluster created after the
emission of an electron due to the photodetachment.
72
6 Results and Discussion
Ag 3d states
∆Ε≈0.6 eV
a)
Intensity (a.u.)
3d3/2
3d5/2
2 nm
4 nm
6 nm
pure Ag
362
364
b)
366
368
370
372
374
376
Binding Energy (eV)
378
380
1,5
368,9
1,4
368,8
368,7
1,3
368,6
368,5
1,2
368,4
Ag bulk
368,3
368,2
FWHM (eV)
Binding Energy (eV)
369,0
1
2
3
4
5
6
Ag nanoparticle size (nm)
7
1,1
Fig 6.24 a) XPS Ag 3d core level of Ag nanoparticles on HOPG as a function of the particle
size. b) Summary of Ag 3d level shifts and broadening of Ag 3d peaks as a function of the
particle size.
6 Results and Discussion
73
The positive core level shifts summarized in fig 6.24 b), were also
observed for metal nanoparticles on other substrates, which were often
interpreted in terms of final state effects [31-36]. One can argue that our results
are due to the insufficient conductivity of HOPG: For deposited nanoparticles
on poorly conducting substrates [34, 35, 105, 133], the number of conduction
electrons and neighboring atoms is limited by the particle size and, hence, the
positive hole, which is the final state of the photoemission process, can be less
efficiently screened. However, if one compares our results to those of Ag
nanoparticles on TiO2 with a band gap bigger that of HOPG [133], it is
questionable why for Ag/TiO2 the core level binding energies can almost
converge to the bulk values, when Ag particle sizes exceeds 3-5 nm in diameter
and about 2 nm in height. However, as shown in Fig 6.24, Ag nanoparticles on
HOPG as large as 4 nm in lateral size and 3 nm in height still show a deviation
of 0.4 eV in the Ag 3d states with respect to the Ag bulk value. Thus, final state
effects cannot totally explain the observed chemical shift, and initial state
effects should also contribute to this observation.
One important initial state effect is the metal-support charge transfer.
Considering an electron charge transfer from Ag nanoparticles to our
supporting material HOPG, one can shed light into the changes in chemical
shifts [141, 142]. However, core level shifts up to 0.6 eV for Ag nanoparticles
on HOPG cannot be completely explained by a charge transfer. In the XPS data
presented in Fig 6.24, Ag nanoparticles, showing core level shifts of 0.2~0.4
eV from bulk value, are thicker than 3 nm. In contrast, metal films, consisting
of 1-3 atomic layers, were modified by the substrate to develop bulk-like core
level binding energies within the film thickness of about 5 ML (1nm) [143].
Furthermore, different work functions between HOPG (4.6 eV) and Ag (~4.3
eV) can likely justify chemical shifts of up to ~0.3 eV, but not shifts of about
~0.6 eV. Thus, even though a considerable amount of Ag to C electron charge
transfer may take place, other mechanisms are necessary to clarify the total
contribution of the Ag 3d core level shift.
Beyond electron charge transfer and screening effects, the shape of the
particles can have an influence on core level shifts. When particles are not
spherical but have a polyhedral form, the density of undercoordinated atoms is
significantly higher. Undercoordinated atoms can alter significantly the
electronic structure of particles by increased lattice strain, resulting in core level
shifts sensitive to the particle geometry [137].
Another initial state effect in metal clusters is the metal-non-metal
74
6 Results and Discussion
transition, which can be responsible for core levels shifts [119, 144]. In cluster
physics, it is well-known that as a metal nanoparticle grows, the energy gaps
between different electronic levels become narrower and above a certain size,
complete metallic bands are developed, resulting in bulk-like electronic
structure [145, 146] (see Fig 6.25 for illustration).
Fig 6.25 Schematic energy level diagram of metal atoms, dimers, clusters (nanoparticles) and
bulk. As the cluster size increases, a continuous metallic band develops [147].
Earlier studies reported about transitions from metallic states of Au
nanoparticles to insulating states with decreasing cluster size [15, 148, 149].
Shifts in the core level binding energy were related to this metal-insulator
transition, but there are still some doubts concerning the origin of these shifts
[144]. In this kind of transition, core level binding energy increases with
decreasing particle size, while valence bands shift to lower energy levels and a
disappearance of states closer to the Fermi level EF takes place at the same time
[150, 151]. Evidence of metal-nonmetal transitions for Ag and Au nanoparticles
were reported when the particle size becomes smaller than 2.3 nm [148, 149].
The particle size regime, at which this transition takes place, is much smaller
than the size regime, where the onset of the core level shifts appears in Ag XPS
data; hence this effect cannot be a dominating factor to explain the Ag 3d core
level shift in our case.
In summary, though TiO2 has lower conductivity than HOPG,
6 Results and Discussion
75
convergence of the Ag 3d state to the bulk state is lower on HOPG compared to
TiO2. This indicates that larger Ag particles on HOPG cannot strongly be
influenced electronically by the substrate via metal-support charge transfer. In
the absence of further information, core level shifts may rather be related to the
geometry (shape) of particles. Metal/nonmetal transitions cannot be responsible
for the observed core level shifts, because these shifts are still observed for
particle sizes (> 3 nm), i.e., for bigger particles than the sizes in which these
transitions usually begin to appear. Further information by means of an Auger
analysis and additional studies on Au nanoparticles grown on HOPG could
enable us to explain the origin of these core level shifts.
Au 4f Core Level Spectra as a Function of Particle Size
In case of Au nanoparticles on HOPG, similar studies were carried out.
Au/HOPG samples, imaged previously by STM, were transferred into the XPS
chamber. Similar fitting as for Ag 3d peaks was carried out for peaks of Au 4f
states. Fig 6.26 a) displays no significant changes in the XPS Au 4f core levels
(about 0.1 eV with respect to the Au bulk value) with decreasing particle size,
even when the particle size becomes smaller than about 5 nm in lateral size and
3 nm in height.
In order to study smaller particle sizes, the sample was sputtered for a
much longer time to increase the defect density, allowing us to detect smaller
Au nanoparticles with higher intensities of the Au 4f peaks. Fig. 6.26 b) shows
changes in the XPS spectra of Au 4f states with decreasing particle size for a
wider particle size regime. In agreement with our spectroscopy data, earlier
results of Au nanoparticles on HOPG have reported no chemical shifts for Au
4f levels (less than 0.1 eV), when particles exceed 1.5 nm in size [152]. It
should be mentioned, that not all the samples were imaged using STM, and
therefore a qualitative XPS estimation for the particle size is assumed in Fig
6.26 b in the case of the smallest particle sizes. This estimation is only reliable
considering a combination of STM and XPS data for those samples already
presented.
76
6 Results and Discussion
Au 4f States
a)
4f 7/2
4f 5/2
~3 nm
Intensity (a.u.)
∆ E ≈ 0 .1 e V
~4 nm
~10 nm
B ulk
80
82
84
86
88
90
Binding Energy (eV)
b)
Au Bulk
~10 nm
~4 nm
~3 nm
~ 1.5 nm
Au 4f
4f7/2
4f5/2
Intensity (a. u.)
∆E ~ 0.1 eV
Decreasing
particle size
< 1.5 nm
80
82
84
86
88
90
92
94
Binding Energy (eV)
Fig 6.26 XPS Au 4f levels as function of particle size a) For the samples already imaged by STM. b) For
the particle size regime from bulk to less than ~ 1.5 nm. The botton 5 Au 4f spectras are presented again
in detail in Fig 6.27.
6 Results and Discussion
77
In Fig 6.27 Au 4f core levels for the smallest Au nanoparticles (< 1.5 nm)
were studied in detail. An inversion of the core level shifts with decreasing
particle size, i.e. positive binding energy shifts of ~0.3 eV with decreasing
particle size, followed by a negative chemical shift (however still positive
respect to Au bulk values), can be observed. Such an change in the direction of
the Au 4f core levels as a function of decreasing particle size has not been
previously found for other studies carried out with Au nanoparticles on
amorphous graphite or HOPG substrates. Different preparation of the substrate
by other groups can justify the different results [139, 152, 153].
Intensity (arb. units)
4f7/2
Au 4f
4f5/2
Decreasing
particle
size
~ 1.5 nm
80
82
84
86
88
90
Binding Energy (eV)
92
Fig 6.27 XPS spectra of Au 4f core levels as a function of particle size. An change in the
direction of the core level shift can be observed in the smallest Au particle size regime (<1.5
nm).
Comparison of Au with Ag XPS Data: Auger Analysis
In Fig 6.24, 6.26 and 6.27, much smaller binding energy shifts of the Au
core levels compared to those of Ag were observed for similar particle size
regimes. In particular, the Au 4f level shifts suffer an inversion for the smallest
particle sizes; first with positive shifts with decreasing particle sizes, followed
by a negative one. In contrast, only positive shifts for Ag nanoparticles with
78
6 Results and Discussion
decreasing particle sizes were observed. These different results for Ag and Au
nanoparticles suggest that these core level shifts cannot be explained within a
simple model, such as the screening effect due to the final state. Other effects
related to the initial state should be taken into account.
An analysis of the kinetic energy shifts of the Auger peaks with
decreasing particle size combined with core level shifts in Ag and Au
nanoparticles can yield information about how to discompose the origin at these
binding energy shifts into initial and final state contributions. For this analysis,
we used an Auger parameter (α) developed by Wagner [106, 154, 155], whose
meaning will be explained. Based on some assumptions, this parameter α will
be deduced.
The core-electron binding energy shift, ∆E, measured in a metal cluster, is
understood as a contribution of both effects (initial and final effects) and can be
written as:
∆E = ∆ε − ∆R + ECoul
(1)
whereas ∆ε represents the change in the initial-state energy, ∆R, the change in
final-state relaxation energy, and Ecoul the coulomb energy due to the unit
charge, left by emission of a photoelectron.
The shift of the kinetic energy by an Auger process, ∆K, is more difficult
to calculate. ∆K is defined as the difference between the energies of the initial
and final state, i.e. before and after the Auger process. At the beginning of an
Auger process, there is a one-hole state, similar to the shift of the core-electron
binding energy (assuming no relaxation before the Auger process). Therefore,
the initial Auger state, EAuger_init, is expressed as in equation (1). The final state
of the Auger process, EAuger_fin, is a two-hole state, with an energy shift given
by:
E Auger _ fin = ∆ε Auger − ∆RAuger + ∆ECorr + 4 ECoul
(2)
In equation (2), terms related to the two-hole states correspond to two
different orbitals and are distinguished by the subscript Auger. ∆εAuger and
∆RAuger are the change in the initial-state energy and the change in the finalstate relaxation energy for these two hole-states by the Auger process,
respectively. ∆ECorr contains the change in the repulsion between the two finalstate holes. The macroscopic coulomb term, 4Ecoul, is 4 times that of the initial
6 Results and Discussion
79
state because two electron charges are missed in the Auger final state and two
positive unit charges are gained. Therefore, the Auger kinetic energy shift, ∆K
can be written as:
∆K = E Auger _ init − E Auger _ fin ⇒
⇒ ∆K = ∆ε − ∆ε Auger + ∆RAuger − ∆R − ∆ECorr − 3ECoul
(3)
By adding equations (1) and (3), one obtains:
∆E + ∆K = 2∆ε − ∆ε Auger − 2∆R + ∆RAuger − ∆ECorr − 2 ECoul
(4)
One can subordinate eq. (4) to a number of restrictive assumptions:
1) For an Auger transition, C´C´´C´´´, C´ is the core level, and C´´C´´´ are the
two Auger final state holes (see Chapter 4, Auger Peaks). It can be assumed that
the initial-state shifts per hole (∆εauger/2) and the change in the initial state
energy, ∆ε, are identical, as long as equal binding energy shifts and identical
relaxation mechanisms for all three core levels take place:
∆ε Auger
2
; ∆ε
(5)
2) The relaxation energy for a two hole state, ∆RAuger is just four times ∆R,
since it contains the interaction of two holes with the twofold screening charge
on an individual atom.
∆RAuger ; 4∆R
(6)
3) Changes in the coulomb correlation energy are significantly of lower order
than other shifts and it can be neglected, whilst no significant differences in the
binding energy shifts for different orbital are assumed.
Taking into account the restrictions 1)-3) and substituting in equation (4):
∆E + ∆K = 2∆R − 2 ECoul
(7)
Note that the Coulomb term “-2ECoul” is of final-state character and cannot be
separated or distinguished from the relaxation term, 2∆R and therefore drops
out in the calculation of ∆E+ ∆K. Thus, eq. (7) can be reduced to:
∆E + ∆K ; 2∆R
(8)
80
6 Results and Discussion
Equation (8) expresses that the sum of the kinetic energy shift of the
Auger Peak and the shift of the core level binding energy equals twice the final
state contribution of the core level shifts.
Auger parameter α is called to the sum of the kinetic energy of the Auger
transition K(C´C´C´´´) from C´, C´, and C´´´ states (see Chapter 4),
respectively, and the core level binding energy E(C´) [154]:
α = BE (C´) + K (C´C´´C´´´)
(9)
Hence, the variation of the Auger parameter, ∆α expressed by eq. (9), leads to:
∆α = ∆BE (C´) + ∆K (C´C´´C´´´)
(10)
If the Auger binding energy is called EAuger, then
∆K (C´C´´C´´´) = −∆Eauger
(11)
and hence;
⇒ ∆α = ∆BE (C´) − ∆Eauger ; 2∆R
(12)
Therefore, the change in the Auger parameter (∆α) is approximately twice the
final state effect. Equation (12) offers a direct way to evaluate the final state
contribution ∆R.
Experimentally, due to the application of Al Kα X ray sources (1486.6 eV),
most Auger transition peaks come from C´VV or C´C´´V excitations28, rather
than C´C´´C´´´ excitation. We used the 3d core level spectra and MVV
Auger spectra for the analysis. In the case of Ag, the strongest electronexcited Auger transition involves a 3d5/2 initial-state hole and two 4d finalstate holes [33]. The top of the 4d band, situated close below the Fermi level,
overlaps the 5s conduction band. For this reason, the Auger transition is of
MVV type. In our case, this analysis could be only carried out for our
Ag/HOPG system, since Au Auger peaks turned out to have too low intensity
due to the small cross section for the Au MVV Auger lines. Applying
formula (9) to our Ag system:
α = E (3d ) + Ek ( MVV )
28
(13)
C and V mean core level and valence level, respectively.
6 Results and Discussion
81
⇒ ∆α = ∆E (3d5/ 2 ) − ∆Eauger (4d ) ; 2∆R
(14)
where M is the 3d5/2 core level orbital and V the 4d valence band orbital.
∆E(3d5/2) is the 3d core level shift and ∆EAuger(4d) the 4d valence level shift.
Using eq. (14), the final state contribution to the binding energy shift of
the 3d core level is ∆α/2 given by:
∆R =
∆α ∆E (3d5/ 2 ) − ∆Eauger (4d )
=
2
2
(15)
∆E(3d5/2) and ∆EAuger(4d) can be directly estimated from our XPS data.
The initial state contribution ∆ε can be obtained by subtracting the final state
contribution ∆R from the total core level shift ∆E(3d5/2):
∆ε = ∆E (3d5/ 2 ) − ∆R
(16)
Many authors have discussed and modified this Wagner’s parameter α
concept [106, 154, 155], concluding that this concept can be only rationalized,
assuming equal binding energies and identical relaxation mechanism for all
three orbital levels. When the two Auger final state holes are in the same shell
as in our case (4d states), involving the valence band, for example MVV Auger
transitions, there may be some limitations. The 4d valence band can be
delocalized and some differences in screening effects between 3d core holes
and 4d valence states may occur. That happens in the same way for Au. Ag has
been studied experimentally and theoretically in terms of its validity with
respect to the Auger parameter approximation [133, 156, 157], pointing out that
in many cases, Auger transition, like Pt MVV, Sn MNN, or Ag MVV for Ag
nanoparticles on Al2O3, amorphous carbon or Ag/TiO2 [133] are valid for this
approximation. In our case, Ag has an intense and sharp Auger MVV (3d-4d)
transition feature in the XPS spectra and parallel shifts of the 3d peak and the
center of the 4d band of Ag have been found, demonstrating the validity of our
Wagner analysis. Evidences that 4d valence holes and 3d core holes of all 4d
metals (except Pd) have the same relaxation in the final state were reported
[156, 157].
82
6 Results and Discussion
Shifts to higher binding energies of the Ag 3d core level and Ag 4d
valence level with decreasing particle sizes are shown in the XPS data in Fig
6.28. As already mentioned, the comparison of both Ag states is necessary in
order to confirm the validation of our Auger analysis.
A g 3 d sta tes
Intensity (a. u.)
3 d 5 /2
3 d 3 /2
B u lk
6 nm
5 nm
4 nm
4 nm
3 nm
~1nm
364
366
368
370
372
374
B in d in g E n e rg y (e V )
376
378
380
Intensity (a. u.)
Ag 4d
Ag bulk
5 nm
x5
4 nm
x 20
0
2
4
6
8
10
Binding Energy (eV)
Fig 6.28 Top: XPS Ag 3d core level as a function of the particle size. Botton: XPS Ag 4d
valence levels as function of the particle size. For each spectrum, different data acquisition
times were used, which are denoted. Note: Low intensities of Ag 4d peaks in comparison with
Ag 3d were observed.
6 Results and Discussion
83
Ag 3d5/2Binding Energy (eV)
The XPS data in Fig. 6.28 are summarized in Fig 6.29, where the binding
energies of Ag 3d core level and Ag 4d valence level were plotted as a function
of the coverage of the deposited Ag metal on HOPG. The coverage was
calculated as a function of the intensity ratio Ag (3d5/2)/C (1s). This comparison
reveals that similar changes of the centers of the 3d and 4d levels take place,
confirming only a deviation of about 0.2. This small deviation justifies the use
of this Auger analysis.
369,2
369,1
a)
369,0
368,9
368,8
368,7
368,6
368,5
368,4
Ag Bulk
368,3
368,2
0
50
100
150
200
250
Ag 4d Binding Energy (eV)
Ag (3d5/2) / C (1s) intensity ratio
6,2
b)
6,0
5,8
5,6
5,4
Ag Bulk
5,2
0
50
100
150
200
Ag (3d5/2) / C (1s) intensity ratio
Fig 6.29 Ag 3d and Ag 4d states as a function of Ag coverage (i.e. particle size) on HOPG. The
changes of the centers of 3d and 4d states confirm a deviation of only ~ 0.2 eV.
84
6 Results and Discussion
Fig 6.30 shows XPS data of Ag MVV Auger lines as a function of Ag
particle size on HOPG used for our Auger analysis. Note that Ag MVV Auger
lines shift to higher binding energies with decreasing particle size.
Auger MVV line
of Ag on HOPG
bulk
Intensity (a.u)
6nm,
5nm,
4nm,
4nm,
3nm,
1nm
1122 1124 1126 1128 1130 1132 1134 1136 1138 1140
Binding Energy (eV)
Fig 6.30 Ag MVV Auger peaks as a function of particle size.
The results of the Auger analysis of the XPS data for Ag on HOPG are
demonstrated in Fig 6.31. The total binding energy shift ∆E, are due to ∆ε- ∆R
(see eq. 16, page 80), in which ∆ε is the initial state contribution and ∆R is the
final state contribution. Both contributions are separately plotted as a function
of Ag particle size. The Auger analysis reveals that the final state effect is about
0.1 eV, independent of the particle size. The initial state contribution
significantly increases with decreasing particle size, resulting to be mostly
responsible for the positive core level shifts of Ag nanoparticles on HOPG. For
a deeper understanding of the different initial state contributions responsible of
these shifts, XPS data of Ag/HOPG and Au/HOPG system will be compared
with the information obtained by the Auger analysis.
6 Results and Discussion
85
Decreasing particle size
Bulk
-0,4
Energy (eV)
-0,2
0,0
0,2
0,4
0,6
final state effect
initial state effect
0,8
-0,1
0,0
0,1
0,2
0,3
0,4
0,5
0,6
0,7
0,8
Ag 3d core level shifts (eV)
Fig 6.31 Final and initial state contribution as a function of Ag 3d core level shifts (i.e particle
size).
Metal-Support Interaction
The Auger analysis shows a minimal contribution of about 0.1 eV for the
final state effect and up to 0.6 eV for the initial state. In contrast, earlier Auger
analysis of Ag nanoparticles on alumina by D. W. Goodman et al [106], showed
a major contribution of the final state effect, in which the initial state effect
contribution was of only 0.1 eV. These different results suggest that the
supporting material can have influence on the electronic structure of Ag
nanoparticles. As already mentioned, the work function of Ag bulk and HOPG
are ~ 4.2 eV and ~ 4.6 eV, respectively, and therefore Ag nanoparticles can
become partially positively charged, contributing to a total positive core level
shift [141, 142]. In the case of Au, the work function is about 5.2 eV, indicating
that Au nanoparticles become negatively charged, causing an opposite effect,
which it is reflected in a negative contribution to the core level shifts. However,
Au nanoparticles still show positive core level shifts with respect to the bulk
value (see Fig 6.26), while electron charge transfer from C to Au tends to
counteract the total positive core level shift with respect to Au bulk values. The
tendency of Au clusters to become more negatively charged than the Ag
86
6 Results and Discussion
clusters can be understood by a higher electron affinity29 of Au clusters
compared with those of Ag clusters [158]. Thus, positive and negative charging
of Ag and Au nanoparticles, respectively, can partially explain the larger
positive chemical shifts of Ag with respect to Au nanoparticles in similar
particle size regime. It should also be mentioned that the inversion from
positive to negative core level shifts with decreasing particle size found for the
smallest Au particles is also in line with the negative charging of Au
nanoparticles. It seems that the negative charging of the Au nanoparticles
partially compensates the positive core level shift caused by other initial states
effects.
Lattice Strain
Rehybridization of the valence orbitals d und sp for Ag (or Au) can result
in an additional positive chemical shift of the core levels [63, 140]. This
happens when the lattice constant increases as a result of a shorter average bond
distance with decreasing particle size [159]. Although the lattice strain can be a
factor to fully understand the larger core level shifts of Ag compared to Au
nanoparticles, due to diverse arguments, which will be exposed, we suggest that
lattice strain is not a dominating factor.
Previous results report that for Ag on carbon almost parallel shifts of the
Ag 3d, 4d, 5sp levels were observed, which do not rationalize the
rehybridization of 4d and 5sp induced by the increased lattice strain for smaller
clusters [33]. Although the measurements of the Fermi level shifts are missing
in our Ag/HOPG system, our results on Ag 3d and 4d shifts data are in
agreement with the data of Wertheim et al. [159]. It is also important to mention
that earlier results suggested that the hybridization of Au clusters is stronger
than those of Cu or Ag, leading to the formation of planar Au cluster up to the
cluster size of ~13 atoms [30]. Since sd hybridization of Au seems to be
stronger than that of Ag, one may expect larger positive core level shifts for Au
than Ag, which is not in agreement with our experimental data.
Role of Undercoordinated Atoms
For Ag and Au nanoparticles on HOPG, no indication of round shapes
was found. Considering that undercoordinated atoms are usually located on
kinks and steps of particles, the concentration of atoms with lower coordination
29
The electron affinity of an element is the energy given off when a neutral atom in the gas
phase gains an extra electron to form a negatively charged ion.
6 Results and Discussion
87
numbers may be higher than for particles with round shapes [62] Previous
results report that undercoordinated atoms can show different electronic
structures than atoms on a perfect surface, or bulk, leading to a chemical shift as
a function of particle size [17]. From the STM data, no significant deviation
between the geometries of Ag and Au nanoparticles on HOPG was found.
Hence, it is not reasonable to assign the difference of the core level shifts of Ag
and Au nanoparticles on HOPG to the particle structure. Therefore, metalsupport charge transfer should be a more important factor in order to explain
this difference.
Metal-Nonmetal Transition
Similar arguments as those given for Ag nanoparticles, can be used for
Au nanoparticles. Metal-nonmetal transition for Au nanoparticles takes place in
a similar Ag particle size regime. However, larger particles than 2-3 nm in
diameter still experiment chemical shifts. Furthermore, parallels 3d, 4d level
shifts in the case of Ag were observed in Fig 6.29, excluding the possibility of
considering a broadening of the electronic structure of the Ag bands, which
would lead to a metal/non-metal transition.
Summary
The Auger analysis to distinguish between final and initial state contributions
shows clearly the importance of the initial state effect (over the final state
effect) in Ag 3d and Au 4d core level binding energy shifts. The observation
that Au 4f level first shifts to higher and then to lower binding energies with
decreasing particle size suggests that this behavior cannot be explained by a
single mechanism. Rather a combination of various mechanisms was suggested.
Lattice strain effects and metal/non-metal transition, together with a final state
effect, cannot justify the different shifts of the Ag and Au core levels, since
these factors should result in similar core level shifts of Ag and Au, or even
larger shifts for Au particles. Our XPS data indicates that the significant
charging of Ag and Au nanoparticles due to metal support interactions can
explain mostly core and valence level shifts of Ag and Au nanoparticles,
together with the inversion observed by the Au core levels. Therefore, Metalsupport interactions play the most important role for the observed changes in
the electronic structure of those particles, previously attributed to geometric
factors. A secondary role plays the shape of particles, in which different
electronic structures of undercoordinated atoms are considered.
88
6 Results and Discussion
W-Oxide Coating on Ag Nanoparticles
Significant difference in the Auger analysis of Ag nanoparticles on
HOPG and Aluminia oxide was found between by Y.D. Kim et. al., [106] and
Goodman et. al., [160]. This difference points out the importance of the
supporting material. For a better understanding of the influence of oxide
substrate, W-oxide was deposited as a layer on Ag nanoparticles on HOPG. Fig
6.32 shows XPS spectra of 3d states of Ag nanoparticles, with a diameter of
about 4 nm, changing with the increase of W-oxide coating thickness. A
decrease of the intensity in the Ag 3d peaks and a shift to lower binding
energies with increasing W-oxide coating thickness is observed. Nevertheless,
Ag 3d core levels of the system W-oxide/Ag/HOPG are shifted to higher
binding energies compared to those of the Ag bulk crystal (368.3 eV).
Therefore, larger positive core level shifts resulting from Ag nanoparticles
without metal-oxide interactions counteract the small negative shift of the Ag
3d states caused by metal-oxide interactions. This negative shift may be
explained by W-oxide to Ag charge transfer and/or an additional screening of
the hole state in Ag by the surrounding W-oxide.
Ag 3d
Ag (4 nm)/ HOPG
Intensity (a.u.)
With increasing
tungsten oxide
coverage
~ 0.2 eV
Ag Bulk
362
364
366
368
370
372
374
376
378
380
Binding Energy (eV)
Fig 6.32 Changes of the Ag 3d core levels with increasing W-oxide coating thickness for Ag
nanoparticles on HOPG with a mean diameter of 4 nm.
6 Results and Discussion
89
The same experiment was carried out with a W-oxide layer growing on
Ag bulk crystal. The XPS data of Fig 6.33 indicates a decrease of the Ag signal
in line with the XPS data of Fig 6.32, however, no chemical shift with
increasing W-oxide coating coverage is observed. It should be mentioned that
significant changes of the Ag particle size while increasing the oxide coating
thickness can be excluded; since intensities of C 1s peak and Ag 3d upon
deposition of W-oxide (coating) were consistent with formation of W-oxide
coating on Ag/HOPG. Thus, it seems that the deposition of oxide layers on
metal-nanoparticle/HOPG could shed some light onto possible charge transfers
between metals and oxides. When metal particles are first deposited on HOPG,
and then different oxide layers are added, interaction of various oxide materials
with metal particles with a similar geometry can be a subject of future study,
which could provide new insight into the interpretation of the core level shifts
data of metal nanoparticles on oxide.
Ag3d
Ag bulk
Intensity (a.u.)
With increasing
tungsten oxide
coverage
362
364
366
368
370
372
374
376
378
380
Binding Energy (eV)
Fig 6.33 Changes of the Ag 3d core levels with increasing W-oxide coating thickness for Ag
bulk.
90
6 Results and Discussion
Intensity (a.u.)
W4f
32
34
36
38
40
Binding Energy (eV)
Fig 6.34 W core levels showing formation of W- oxide coating. The W 4f levels are shifted to
higher binding energies compared to those of metallic W, indicating W-oxide formation.
6 Results and Discussion
91
6.2 Au Nanoparticles on SiO2/Si
Au nanoparticles with different sizes supported by silica were prepared
and imaged by STM. To shed light onto the electronic structures of the small
Au nanoparticles on silica, XPS measurements were performed. These data
were compared with our results of Au nanoparticles on HOPG and those of
previous studies on Au nanoparticles formed on Titania.
STM Analysis
STM images were taken to estimate the Au particle size on Silica (Fig
6.35-6.37). In contrast to HOPG, the Silica substrate was not sputtered for these
3 different samples. The high stability of Au nanoparticles on Silica enable to
prepare narrow particle size distributions without sintering effects, as it has
been confirmed with our STM data. With increasing Au coverage, a three
dimensional growth was observed, where the average particle diameter and
height increases.
•Deposition rate: ~ 4 ML/min
•Coverage: ~4 %, 4 ML
•Particle density: ~ 104 µm-2
25
40
35
Count
Count
25
20
15
10
Height: ~ 0.7 nm
20
diameter: ~ 1.4 nm
30
15
10
5
5
0
0
1
2
3
4
5
6
7
8
Particle diameter (nm)
9
10
0
0
1
2
3 4 5 6 7 8
Particle height (nm)
9 10
Fig 6.35 STM image of Au nanoparticles on native silica layers on Si wafer (193 x 193 nm2). Tunneling
parameters: 2V, 0.1 nA. Lower pictures: Diameter and height distribution of the particles.
92
6 Results and Discussion
For the determination of the particle diameter, one should note that the
apparent particle size is generally significantly larger than the real size due to
the limited resolution of STM tips. As already mentioned, an overestimation of
the particle diameter in STM should be assumed [103, 161]. The particle height
can be more accurately determined, even though some error bar should be
considered (several tens of nanometers), in particular when a high sample bias
is used for STM imaging [139]. The emission current was maintained constant
at a value of 17 mA for all the samples. Fig 6.35 shows a STM image of Au
nanoparticles grown on Silica prepared by evaporating Au for 1 minute at room
temperature. In the cases of Fig 6.36 and Fig 6.37, Au was evaporated for 5 and
25 minutes, respectively.
•Deposition rate: ~ 1 ML/min
•Coverage: ~10 %, 5 ML
•Particle density: ~ 104 µm-2
40
25
35
Diameter: ~3 nm
30
Count
25
Count
Height: ~ 1nm
20
20
15
15
10
10
5
5
0
-1
0
1
2
3
4
5
6
7
Particle diameter (nm)
8
9
10
0
0
1
2
3 4 5 6 7 8
Particle height (nm)
9 10
Fig 6.36 STM image of Au nanoparticles on native silica layers on Si wafer (124 x 124 nm2). Tunneling
parameters: 4.7 V, 0.1 nA. Lower pictures: Diameter and height distribution of the particles.
6 Results and Discussion
93
•Deposition rate: ~ 0.8 ML/min
•Coverage: ~24 %, 22 ML
•Particle density: ~ 104 µm-2
20
32
Height: ~ 4.5 nm
28
16
Diameter: ~<10 nm
Count
24
Count
20
12
16
8
12
8
4
4
0
0
1
2
3
4
5
6
7
8
Diameter (nm)
9
10 11 12
0
0
1
2
3 4 5 6 7 8
Particle height (nm)
9 10
Fig 6.37 STM image of Au nanoparticles on native silica layers on Si wafers (123 x123 nm2).
Tunnelling parameters: 4.7 V, 0.1 nA. Lower pictures: Diameter and height distribution of the
particles.
In contrast to Au nanoparticles deposited on HOPG, no defined polygonal
structure in the geometry of the particles could be clearly identified. In order to
image these samples, high bias voltages (~ 2V- 4.5V) were employed. High
voltages are necessary due to the much lower conductive of silica compared to
HOPG. Even though high bias voltages usually induce some artificial effects,
no sintering or tip effects were observed. This may indicate that a strong
interaction between Au nanoparticles and the silica substrate should take place.
This interaction may be dominated by significant charge transfer effects,
already found in the case of HOPG. Further information about this kind of
interaction is explained in the analysis of our XPD data in the next section.
Fig 6.38 displays the growth mode of Au on silica, in which particles
undergo a three dimensional growth with increasing particle size. An estimation
of the parameters used for the preparation of the Au nanoparticles of Figures
6.35-6.37 is presented in Table 3.
6 Results and Discussion
Particle Height (nm)
94
11
10
9
8
7
6
5
4
3
2
1
0
Au Growth Mode
0
1
2
3
4
5
6
7
8
9
10 11
Particle diameter (nm)
Fig 6. 38 Particle diameter versus particle height of Ag nanoparticles grown on Silica.
Deposition
Rate
(ML/min)
Particle
Density
Au
Diameter
Height
Coverage
(ML, %)
Fig 6.35
~1.4 nm
~ 0.7 nm
4 ML, 4 %
4 min-1
~ 104 µm-2
Fig 6.36
3 nm
1 nm
5 ML, 10 %
1 min-1
~ 104 µm-2
Fig 6.37
< 10 nm
~ 4.5 nm
22 ML, 24 %
0.8 min-1
~ 104 µm-2
Table 3 Estimation of parameters used for the preparation of Au nanoparticles grown on Silica.
6 Results and Discussion
95
XPS Results
A XPS study of different samples was carried out. In this study, the
samples imaged by STM were included. In Fig 6.39, core level shifts of the Au
4f states as a function of Au coverage are displayed. For (c), (f) and (g) the
average particle heights (widths) estimated on the basis of the STM data in Fig
6.35-6.37 are approximately 0.7 (<1.4) nm, 1 (3) nm and 4.5 (<10) nm,
respectively. With decreasing particle size, positive core level shifts of the Au
4f doublet features are observed. This chemical shift is in line with previous
results, in which metal nanoparticles on supporting materials with relatively
poor conductivities, as alumina and Titania, also show this behavior [31, 33,
106, 134, 136, 162-164]. A peak deriving from the Si substrate (a satellite
feature of Si) disappears besides the Au 4f peak with increasing of Au
coverage.
Intensity (arb. units)
Au 4f
Si
~ 1 eV
increasing
particles
size(diameter)
78
h (Bulk)
10 scan
g (~10 nm)
f (~3 nm)
e
d
c (~ 1.4 nm)
b
a
10 scan
80
82
15 scan
20 scan
20 scan
25 scan
30 scan
50 scan
84
86
88
90
92
94
Binding Energy (eV)
Fig 6.39 XPS Au 4f levels for Au nanoparticles on SiO2/Si. From a to h, Au coverage increases.
A Si satellite peak with decreasing particle size appear. For each spectrum, different data
acquisition times were used, which are denoted.
96
6 Results and Discussion
Fig 6.40 shows the core level shifts of the Au 4f states as a function of
amount of deposited Au. A maximum core level shift of about 1 eV with
respect to the values of Au bulk crystals (84.0 eV) is found. It is important to
mention that the lowest Au coverage, which has been prepared, is estimated to
be below 10 % of a monolayer. Comparing these results with an equivalent Au
coverage on HOPG (see Chapter 6.1), much less core level shifts were found.
An explanation can be the higher conductivity of substrate HOPG with respect
to SiO2/Si substrate. A higher C to Au electron charge transfer can very
efficiently screen much hole state compared to the case of Au/SiO2, justifying
these different results. In the past, core level shifts of ~ 1.6 eV were observed
for similar Au coverages on silica thin films on molybdenum surfaces
(SiO2/Mo(110) [134, 165, 166]. These shifts are larger than those found in our
system Au/SiO2/Si with a similar Au coverage regime. Our data are rather
analogous to the core level shifts found for Au on Titania [165, 166].
4 f level Binding Energy (eV)
Au 4f level shift
85
Au Bulk
84
0
2000
4000
Au evaporation time (second)
Fig 6.40 Centers of Au 4f7/2 peaks as a function of amount of Au deposite (i.e. particle size)
To shed light onto the origin of these different results between our system
Au/SiO2/Si and the system Au/SiO2/Mo(110) [134], Si 2p levels were collected
and summarized as a function of particle size as shown in Fig 6.41. In case of
Au/SiO2/Mo(110), core level shifts of up to 1 eV were observed in the Si 2p
level with increasing Au coverage for the system Au [134]. In contrast, Fig 6.41
shows that a Si 2p level shift of only 0.1 eV could be found.
6 Results and Discussion
97
a)
Intensity (arb. units)
S i 2 p S ta te s
In cre a sin g
p a rticle size
h
g
f
e
d
c
b
a
(<10 nm )
(3 nm )
(1.5nm )
~ 0.1 eV
96
98
100
102
104
106
B inding E nergy (eV )
b)
Binding Energy (eV)
94,5
Si 2p level shift
95,0
95,5
96,0
0
1000
2000
3000
4000
Au evaporation time (seconds)
Fig 6.41 a) XPS spectra of Si 2p level as a function of Au coverage. b) Centers of the Si 2p
peaks as a function of amount of Au deposited.
This result can be understood by taking different thickness of the silica
layers used for these two systems into account. For the system
Au/SiO2/Mo(110), silica layers of about 3 nm in thickness were used; whereas
the native silica films used for our substrate SiO2/Si were thinner than 1 nm
(estimated by the intensities of the Si 2p peaks of Si4+ and Si0). A thicker oxide
film can suffer from the charging problem, i.e. the positive hole state in a
thicker insulating film can less efficiently be screened compared to the case of a
98
6 Results and Discussion
thinner film, causing a more pronounced core level shift for the thicker film as a
function of metal ad-layer coverage. It is worth to mention that core level shifts
with decreasing particle size of Au nanoparticle on silica observed by Boyen et
al. [167], are in good agreement with our data.Furthermore, our XPS data point
out to dissimilar valence band structures of Au nanoparticles compared to the
bulk electronic structure.
Previous studies, theoretically as well as experimentally, predict this kind
of redistribution of the valence band structure with decreasing particle size
[104, 162, 168]. They suggest that this redistribution of the valence band is due
to an increase in the number of undercoordinated atoms with decreasing particle
size. In Fig 6.42, XPS valence band spectra taken from some of the same
samples of Fig 6.39 are displayed. Two different Au-induced peaks can be
found centered at ~4, and ~6.5 eV, which are denoted as peak A (5d5/2) and B
(5d3/2), respectively. The peak A shifts to higher binding energy with decreasing
particle size, which was also found for the Au 4f level. In contrast to the
behavior of the peak A, peak B does not undergo any shift upon decreasing
particle size. The d-band splitting for the largest Au nanoparticles is 2.6 eV,
which is close to the bulk value (2.7 eV). The splitting becomes narrower with
decreasing particle size, showing 1.9 eV of the splitting for the Au
nanoparticles of about 2 nm in diameter (estimated by the Au 4f level shift). A
similar narrowing of the splitting between 5d5/2 and 5d3/2 states was found
previously for Au/Cu alloy systems [169].
The apparent d-band splitting term can be described by convolution of the
atomic spin-orbit splitting and electronic band term. As the number of directly
neighboring Au atoms decreases, the electronic band term is decreasing,
resulting in a decrease of the 5d-band splitting [169, 170]. In the Au/Cu alloy
system, the 5d-splitting was found to correlate with the dilution of Au [169]. As
the number of directly neighboring Au atoms decreases with decreasing particle
size, the electronic band term contributing to the 5d-band splitting becomes
smaller, resulting in a decrease of the d-band splitting. A similar decrease of the
d-band splitting as a function of particle size was found for Au nanoparticle on
Zirconia [162].
Besides the band splitting, the band narrowing is also obvious with
decreasing particle size, which is in line with previous theoretical studies,
suggesting a different valence d-band structure with increasing number of
undercoordinated atoms [17].
6 Results and Discussion
99
Particle size (diameter)
Intensity (arb. units)
h (bulk)
g (~10nm)
f (~3nm)
e (<3nm)
d (>1.4nm)
A(Au 5d5/2)
2.6 eV
A
B
B(Au 5d3/2)
1.9 eV
-2 -1 0
1
2
3
4
5
6
7
8
9 10
Binding Energy (eV)
Fig 6.42 Valence band structures as a function of Au coverage. A spectrum of a bare Si wafer was
subtracted from each spectrum. The five spectra were taken from the samples of the spectra h-d in Fig
6.38. One can observe the band narrowing and the decrease of the splitting between the peaks A and B.
Summary
Different electronic properties for Au nanoparticles formed on HOPG and
silica were found. It was previously shown that Au nanoparticles on HOPG
should be partially negatively charged. In contrast, Au nanoparticles on silica
result in very weak metal-support interaction. The electronic structure of Au
nanoparticles on silica was compared to those on other substrates such as
Zirconia and Titania to shed light onto the metal support interaction. Positive
core level shifts compared to the bulk Au, mostly attributed to the final state
100
6 Results and Discussion
effects, were found. Moreover, a decrease of the average coordination
number of Au nanoparticles was shown to significantly change the electronic
structure of Au nanoparticles. This may be responsible for interesting
chemical properties of this particle as catalyst founded in further
experiments. Further studies about oxidation and reduction behaviors of Au
nanoparticles on HOPG and silica have shown the importance of these
electronic properties (see Chapter 7).
6 Results and Discussion
101
6.3 Bimetallic (Ag-Au) Nanoparticles grown on HOPG
Bimetallic (Ag-Au) nanoparticles supported by HOPG were prepared
with different sizes and imaged by STM. To shed light onto the electronic
structures of these nanoparticles, XPS measurements were performed. This data
was compared to our results of pure Ag nanoparticles grown on HOPG. The
composition of the bimetallic particles regarding whether the particles mix as
alloy type (miscible system) or as core-shell type (phase separated) is discussed
by means of an XPS analysis.
STM Analysis
STM images were taken to estimate the size and growth mode of
bimetallic (Ag-Au) particles on HOPG (Fig 6.43-6.45) under UHV conditions
at room temperature.
•Defect density: ~4 % of a ML
• Particle coverage: ~7 %, 4 ML
•Particle density: 104 µm-2
45
25
40
Height: ~ 0.8 nm
30
Counts
Counts
35
Diameter: ~ 2.5 nm
20
15
10
25
20
15
10
5
5
0
0
1
2
3
4
5
6
7
8
9 10 11 12
Particle diameter (nm)
0
0
1
2
3
4
5
6
7
8
9 10 11 12
particle height (nm)
Fig 6.43 STM image of bimetallic nanoparticles (Ag@Au)30 grown on HOPG (223.6 x 223.6
nm2). Tunnelling parameters: 1.78 V, 0.15 nA. Lower picture: Size distribution of the particles.
30
Ag@Au means that first Au was evaporated on HOPG and subsequently Ag.
102
6 Results and Discussion
All samples were prepared by sputtering HOPG between 8-10 seconds at
an Ar pressure of 5 x 10-5 mbar. In the case of samples of Fig 6.43 and Fig
6.44, Au was first evaporated for 16 minutes, and subsequently, Ag was
evaporated for 20 and ~ 7 seconds on each sample, respectively. Sample of Fig
6.45 was prepared by evaporating for 30 seconds of Ag first, and then, 4
minutes of Au. The emission current during evaporation of Ag, or Au, was
maintained constant for the three samples, at a value of 15 mA.
•Defect density: ~4 % of a ML
•Particle Coverage: ~19 %, 5 ML
•Particle density: ~ 104 µm-2
20
40
Diameter: ~ 3.5 nm
35
counts
counts
Height: ~ 1nm
30
15
10
25
20
15
10
5
5
0
0
0
1
2
3
4
5
6
7
8
9
particle diameter (nm)
10 11 12
0
1
2
3
4
5
6
7
8
9 10 11 12
particle height (nm)
Fig 6.44 STM image of bimetallic nanoparticles (Ag@Au) grown on HOPG (194.2 x 194.2 nm2).
Tunnelling parameters: -1.89 V, 0.045 nA. Lower picture: Diameter and height distribution of particles.
Narrow size particle distributions were obtained for all the samples. With
increasing of particle size, an increase in height was also observed. This
suggests a three dimensional growth pattern, which is in line with previous
results about the growth mode observed by pure Ag (or Au) nanoparticles on
HOPG (See Chapter 6.1). A better resolution is necessary in order to identify a
possible polyhedral structure of these particles.
6 Results and Discussion
103
•Coverage: ~23 %, 7 ML
•Particle density: ~104 µm-2
•Defect density: ~4 % of a ML
60
Diameter: ~4.5 nm
25
50
Height: ~ 1.5 nm
40
Counts
counts
20
15
10
5
30
20
10
0
0
1
2
3
4
5
6
7
8
9 10 11 12
particle diameter (nm)
0
0
1
2
3
4
5
6
7
8
9 10 11 12
particle height (nm)
Fig 6.45 Two STM images of bimetallic nanoparticles (Au@Ag) on HOPG. Left: 518 x 518 nm2.
Right: 225 x 225 nm2. Tunnelling parameters: -1.97 V, 0.08nA. Lower picture: Diameter and
height distribution of the particles.
Evidence was found that bimetallic (Ag-Au) particles agglomerate on
non-sputtered HOPG in a different way than Ag (or Au) nanoparticles. Ag
nanoparticles on non-sputtered HOPG tend to form smaller particles than Ag
particles growing on sputtered HOPG (see Chapter 6.1). This change in particle
size depending on whether the substrate is sputtered or not, was also found for
bimetallic (Ag-Au) nanoparticles. Furthermore, bimetallic particles seem to
interact with each other in a different manner, building special nanostructures.
Fig 6.46 shows two STM images of the same sample in which diffusion and
agglomeration of single bimetallic (Ag-Au) particles forming larger islands
with ramified shapes can be observed. For the preparation, Au was evaporated
for 9 minutes and then Ag for 4 minutes by using a constant emission current of
17.5 mA. These special nanostructures suggest that the growth kinetic of
bimetallic (Ag-Au) metal particles on non-sputtered HOPG surfaces may differ
from that of monometallic particles. Another area of this sample is shown in Fig
104
6 Results and Discussion
6.47. In contrast to Fig 6.46, this area was slightly sputtered during 3 seconds at
an Ar pressure of 5 x 10-5 mbar. Taking into account the particle size
distribution here, it can be assumed that islands (~ 30 x 10 nm2), observed in
Fig 6.46, are composed of several single bimetallic particles ( ~ 10-12 nm2).
Fig 6.46 Two STM images of the same sample of bimetallic nanoparticles (Ag@Au) grown on
non-sputtered HOPG. Left) 500 x 500 nm2; STM parameters: -3.7 V, 0.88 nA. Right) 233.6 x
233.6 nm2; STM parameters: -3.7 V, 0.33 nA. Note: Especial structures are formed, compound
of several single particles.
Note that for each STM image, the mean particle size was calculated using the
half maximum of the profile of the particle. An overestimation of the size due
to convolution effects of the tip is assumed using this method (see Chapter 6.1).
The estimation of the coverage area was done comparing the STM data with a
estimation based upon the intensities ratios I[Ag(3d)/C(1s)] and I[Au(4f)/C(1s)]
by means of XPS. Using this estimation, the half maximum of the particle
profile was multiplied by a correction factor of 60 % as explained in Chapter
6.1 (see Chapter 6.1: “Combined STM/XPS quantitative analysis of the lateral
particles size”)
In the next section, electronic structures between bimetallic (Ag-Au)
nanoparticles and Ag nanoparticles as a function of the particle size are
compared. For this purpose, mean particle sizes were corrected in both systems.
Moreover, it should be mentioned that only an error bar of approx. several tens
of nanometer affects particle height.
6 Results and Discussion
105
14
30
12
20
8
counts
counts
Height: ~ 3.7 nm
25
Diameter: ~ 10-12 nm
10
6
4
2
15
10
5
0
0
2
4
6
8
10
12
14
16
18
0
20
0
particle diameter (nm)
2
4
6
8
10
12
14
16
18
20
particle height (nm)
Fig 6.47 STM image of bimetallic nanoparticles (Ag@Au) grown on slightly sputtered HOPG:
491.2 x 320 nm2. Tunneling parameters: - 3.7 V, 0.8 nA. Lower picture: Diameter and height
distribution of the particles. Broader particle size distribution was found compared with
samples of Fig 6.43-6.45.
5,0
4,5
Alloy Particles
Growth Mode
4,0
Height (nm)
3,5
3,0
2,5
2,0
1,5
1,0
0,5
0,0
0,0
0,5
1,0
1,5
2,0
2,5
3,0
3,5
4,0
4,5
5,0
Diameter (nm)
Fig 6. 48 Particle diameter versus particle height of Alloy nanoparticles grown on sputtered
HOPG.
106
6 Results and Discussion
Bimetallic
particles
(Ag/Au)
Diameter
Height
Particle Coverage
(ML, %)
Defect
Density
(% of a ML)
Particle
Density
Fig 6.43
~1.4 nm
~ 0.7 nm
4 ML, ~7 %
~4 %
~ 104 µm-2
Fig 6.44
3 nm
1 nm
5 ML, ~19 %
~4 %
~ 104 µm-2
Fig 6.45
< 10 nm
~ 4.5 nm
7 ML,~23 %
~4 %
~ 104 µm-2
Table 4 Estimation of parameters used for the preparation of Alloy nanoparticles grown on
sputtered HOPG.
6 Results and Discussion
107
XPS Studies
Fig 6.48 shows 4 different XPS spectra of the Ag 3d core levels for the
samples in Fig 6.43-6.45. Each peak was fitted by using a combination of 2
Gaussian functions. A main peak and a shoulder at higher binding energies
have been combined to determine the binding energies of the Ag 3d states
accurately. Positive core level shifts in the Ag 3d states of only about ~ 0.1-0.2
eV with respect to the pure Ag bulk value (368.3 eV) were found. Ag 3d core
level shifts of bimetallic particles are summarized (~ 0.2 eV) and compared
with those shifts observed in pure Ag nanoparticles on HOPG (~ 0.4–0.6 eV) in
similar particle size regimes (see Chapter: 6.1 XPS analysis). Thus, much less
Ag 3d level shifts were found in the case of bimetallic particles. This implies
that Au atoms from Ag-Au nanoparticles may influence the electronic structure
of the Ag atoms. Previous results about the electronic properties of bimetallic
systems by W. Goodman and J. Rodriguez have reported about a perturbation
on the electronic structure of metal atoms supported on well-defined surfaces
such as Ni, Cu and Pd [171]. This influence was found to affect the electronic
structure of both constituents of those bimetallic systems.
In our system (Ag-Au)/HOPG, It is suggested that this influence can be
more important if particles form highly miscible alloy systems instead of
forming core-shell structures, such as AucoreAgshell or AgcoreAushell. In the case of
alloy formation, mixing between different core and shell atoms should takes
place, resulting in an interaction between shell and core atoms. Obviously, this
interaction is expected to be stronger than in case of core-shell formation.
Different Ag 3d level shifts in similar particle size were found when one
compares bimetallic particles with monometallic Ag particles. This difference
in the electronic structure of the bimetallic particles with respect to the single
Ag nanoparticles can be attributed to the Au atoms, which may perturb the
electronic structure of the Ag atoms in the bimetallic particle. In the past, alloy
formation in bimetallic particles was predicted theoretically [172] and found
experimentally [173] even for metals known to be immiscible in the bulk phase
(e.g., Cu-Ag, Fe-Ag). This usually happened after reducing their particle size to
less than approx. 2-3 nm. It is reasonable to think that changes in the electronic
structure are due to the diffusion and subsequent effect of mixing of dissimilar
core and shell atoms (i.e. alloy formation). If particles form miscible alloy
systems, the positive hole created during photoemission in Ag atoms could be
more efficiently screened by the surrounding Au atoms inside the alloy particle.
108
6 Results and Discussion
Ag 3d states
1.5 nm
Intensity (a.u.)
~ 0.1 eV
2 nm
~ 0.2 eV
3 nm
pure Ag
362 364 366 368 370 372 374 376 378 380
Binding Energy (eV)
Binding Energy (eV)
369,0
Ag nanoparticles
Ag-Au bimetallic
nanoparticles
368,9
368,8
368,7
368,6
368,5
368,4
Ag bulk
368,3
368,2
1
2
3
4
5
nanoparticle size (nm)
6
7
Fig 6.49 a) XPS spectra of Ag 3d states of (Ag-Au) nanoparticles as a function of particle size
and. b) Summary of the Ag 3d level shifts of Ag nanoparticles and bimetallic (Ag-Au) particles
as a function of particle size.
6 Results and Discussion
109
Moreover, metal-support charge transfer, between Au (or Ag)
nanoparticles and the supporting material HOPG was reported previously in
Chapter 6.1. As a consequence of this mechanism, Au and Ag atoms become
partly negatively and positively charged, respectively. This mechanism could
lead to charge transfer processes (probably not only between support and
particle, but also between Ag and Au particles), which can party compensate
Ag 3d level shifts due to the final state (positive hole) in Ag particles.
In addition, quantum-size effects in the case of pure Au (or Ag)
nanoparticles are well known in this particle size regime (~ 1-3 nm), leading to
metal-nonmetal transitions [174]. It is likely that this kind of transition also
takes place in Ag-Au bimetallic nanoparticles, affecting to the composition of
the particle, or changes in different lattice strains, etc [167, 171, 175, 176]. It is
probable that the role of undercoodinated Au and Ag atoms is different;
inducing some changes in the electronic structure.
Summary
Ag 3d binding energies (in the same particle size regime) were compared
between bimetallic (Ag-Au) nanoparticles and pure Ag nanoparticles of the
same size regime grown on sputtered HOPG. Based on this, alloy formation
instead of surface-segregated (i.e. core-shell formation) was suggested. A
combination of metal-support interaction and strong Ag-Au interaction due to
intermixing atomic processes of Ag and Au atoms, could explain this. These
electronic changes can dramatically alter the chemical and catalytic properties
of the constituents of the particle. For this reason, alloy metal particles can be
promising candidates as catalysts for a multitude of reactions.
110
7 Chemical Behaviours of Ag and Au
Nanoparticles
The information in this PhD-thesis about the electronic and geometric
structure of Ag and Au nanoparticles grown on HOPG and SiO2/Si substrates
was employed in further studies about chemical properties of these particles by
D. C. Lim et al [177]. Size-dependent chemical behavior of these particles was
extensively investigated in our group by D. C. Lim et. al., and reported
somewhere else [54, 103, 104]. Most important results about oxygenchemisorption and CO-oxidation experiments of Ag and Au nanoparticles on
HOPG and silica are summarized here. To study the chemical behavior, these
supported nanoparticles were oxidized with atomic oxygen and subsequent COoxidation experiments were performed. In the cases of Ag nanoparticles on
HOPG, it was found that the oxygen uptake of smaller Ag nanoparticle is
higher than that of larger particle and bulk-like. Depending on the particle size,
different oxygen species were identified.
Intesnsity (arb. units)
O 1s state
Dissolved oxygen
Active Oxygen
species
(Ag2O/AgO)
before CO exposure
after CO exposure
Oxygen_HOPG
526
528
530
532
534
536
538
Binding Energy (eV)
Fig 7.1 XPS spectra of O 1s states for Ag nanoparticles on HOPG (~ 6nm) after exposition to
atomic oxygen and subsequent reduction by CO at room temperature. By D. C. Lim et al.,
[103].
111
112
7 Chemical Behaviours of Ag and Au Nanoparticles
In case of Ag nanoparticles on HOPG larger than 3 nm in diameter, two
different oxygen species were observed. One of these species (Ag2O and/or
AgO) reacts with CO to form CO2 at room temperature. The other one (AgOx),
which is inert towards CO-oxidation, is attributed to dissolved oxygen in Ag
(See Fig 7.1). Only this non-reactive oxygen species was found for the smaller
particles (< 3). This oxidation-reduction cycle observed for particles bigger than
3 nm, was reproduced several times (i.e, reversible) and can be observed by Ag
3d core level shifts spectra. (see Fig 7.2). Details of this size-selectivity in the
chemical behavior are described by D. C. Lim et. al.[104].
Intensity (arb. units)
Ag 3d state
O exposure
O exposure
CO 1000L
O exposure
CO 1000L
O exposure
CO 1000L
O exposure
CO 1000L
O exposure
366
372
378
Binding Energy (eV)
384
Fig 7.2 XPS spectra of Ag 3d level for Ag nanoparticle (~6nm) after exposing to atomic
oxygen and subsequent reduction by CO at room temperature. The oxidation and reduction
cycle is reversible. By D. C. Lim et al., [103].
Similar experiments with Au nanoparticles on HOPG were carried out.
For Au nanoparticles on HOPG smaller than ~10 nm, two different oxygen
species were detected: Au-oxide (Au2O3), which is catalytically active towards
CO-oxidation and subsurface oxygen, which does not react with CO (see Fig
7.3 and Fig 7.4). For Au film, in contrast, only single oxygen species could be
7 Chemical Behaviours of Ag and Au Nanoparticles
113
Intensity (arb. units)
identified, which can be attributed to Au-oxide. In case of Au nanoparticles
smaller than 3 nm, the catalytically active species (Au2O3) was not detected.
O 1s States
Active species:
Au-oxide
subsurface oxygen
O exposure
CO 1,000 L
CO 3,000 L
526
as prepared Au
528
530
532
534
536
Binding energy (eV)
Intensity (arb. units)
Fig 7.3 XPS spectra of O 1s states of Au nanoparticles with a mean size of 5 nm after
expostion to atomic oxygen and subsequent reduction by CO at room temperature. Active and
inert species can be discriminated. By D. C. Lim et al., [104]
Au 4f
CO 3,000 L
CO 1,000 L
O exposure
as prepared Au
82
84
86
88
90
92
94
Binding energy (eV)
Fig 7.4 Au 4f states for Au nanoparticles with a mean size of 5nm after exposition to atomic
oxygen and subsequent reduction by CO at room temperature. By D. C. Lim et al., [104].
114
7 Chemical Behaviours of Ag and Au Nanoparticles
For Au nanoparticles with different sizes (approx. between 1 and 10 nm)
grown on silica, different oxidation patterns were found depending on the
particle size [54]. For Au particles between approx. ~1.4 and ~ 3 nm in width,
oxidation and subsequently a reduction of the oxidized particles with CO was
observed (see Fig 7.5). This behavior was found to be reproducible. In case of
particles smaller than ~1.4 m, the catalytically active species (Au2O3) was not
detected anymore, indicating that these particles in this size regime have a
similar behavior to those of the smallest Au nanoparticles (< 3nm) on HOPG.
Intensity (arb. units)
Different types of oxygen species formed on Au metal nanoparticles as a
function of particle size can explain this catalytic behavior.
as prepared
after atomic oxygen
92
90
88
86
84
82
after atomic oxygen
after CO
80
92
90
Binding Energy (eV)
88
86
84
82
80
Fig 7.5 XPS spectra of Au 4f levels of Au nanoparticles (between ~1.4 and ~3 nm size) grown
on SiO2/Si. Au nanoparticles were exposed to atomic oxygen and subsequently reduced by CO.
This cycle is reproducible. By D. C. Lim et. al.,[54].
Summary
Oxidation and CO-oxidation reactivity of Ag and Au particles was found to
be very sensitive to the particle size. This is likely due to the different oxide
species formed in the nanoparticles depending on size. Among them, novel
oxidation species were found on Ag and Au nanoparticles, which are absent
in bulk crystal. These species play a passive role for the room temperature
CO-oxidation. Nevertheless, to understand the structures of catalytically
active Ag and Au nanoparticles on different substrates, these new oxygen
species have been studied elsewhere by our group [54, 103, 104], suggesting
to have an important role in the size selectivity of heterogeneous catalysis.
8 Conclusion
Electronic and geometric properties of Ag and Au nanoparticles grown on
various substrates by means of evaporating metal atoms were investigated at
room temperature under ultra high vacuum conditions. X-ray Photoelectron
Spectroscopy and Scanning Tunnelling Microscopy were used for this purpose.
Narrow particle size distributions were prepared on Highly Ordered Pyrolytic
Graphite (HOPG) and silica (SiO2/Si) surfaces, which enable us to study in this
regime particle size effects. The motivation relies on term of fundamental
science as well as on the fact that nanoparticles are potential candidates for the
heterogeneous catalysis of various reactions.
To answer the question how the nature of the substrate can alter the
intrinsic properties of the particles, experimental results of Ag and Au
nanoparticles on HOPG and amorphous silica layer were compared. First, the
system Ag/HOPG (Ag nanoparticles grown on HOPG) was investigated for
different surface treatments. When the HOPG surface was sputtered by Ar-ion
bombardment, much higher Ag particle densities could be observed, implying
that the defect sites play an important role for the nucleation of the
nanoparticles. The electronic and geometric properties obtained for this system
were compared to the system Au/HOPG. Another object of study was the
system Au/SiO2/Si. The high density of defects of the amorphous silica layer
increases stability of the particles. Finally, bimetallic Ag-Au particles with
uniform sizes were prepared on HOPG and investigated. In all measurements,
STM and subsequently XPS measurements were performed. The most
important results on these systems are the following:
Ag on HOPG: Different growth behaviors compared with other Van der
Waals substrates (WSe2) were found, indicating subtle changes in the metalsupport interaction between substrates. Changes in the electronic structure of
the Ag particles with decreasing particle size, in term of Ag 3d core level shifts,
were observed. These shifts are related to initial state effects as well as final
state effects. Core level shifts of Ag nanoparticles on HOPG can be influenced
not only by particle size but also by other factors such as particle geometry and
substrate. In addition, evidence of different Ag-W-oxide interactions was found
for Ag particles compared to Ag bulk crystal.
115
116
8 Conclusion
Au on HOPG: No significant differences of the particle structure and
growth between Au and Ag particles were found. Three dimensional growth
and hexagonal shapes with the existence of many steps between different
terraces could be identified. This information suggests that particle growth is
kinetically limited at room temperature. Different electronic behaviors between
Au and Ag nanoparticles on HOPG were found. XPS studies on Au/HOPG
confirmed that initial state effects are important for the core level shifts. Metalsupport interaction and different electronic structure of undercoordinated atoms
can be responsible for this. Comparison of Ag and Au on HOPG by means of
Auger analysis suggests that Ag and Au nanoparticles are positively or
negatively charged, respectively, due to charge transfers between substrate and
metal particles.
Au on SiO2/Si: Au particles were found to grow in a three dimensional
mode on SiO2/Si. The electronic properties of Au nanoparticles smaller than
~10 nm on this substrate were shown to be significantly different from the bulk
structure. Au 4f core level shifts are larger than in case of HOPG, but similar to
core level shifts observed for Au nanoparticles on poorly conducting surfaces
(e.g. Alumini and Titania). Furthermore, it was observed that a decrease in the
average coordination number of Au nanoparticles on SiO2/Si significantly
perturb the electronic structure of the valence band of the particles.
Bimetallic (Ag-Au) Nanoparticles on HOPG: It was demonstrated that
bimetallic Ag-Au nanoparticles can be prepared on HOPG with relatively
narrow size distributions. XPS studies suggest that the Ag 3d states in these
particles can be influenced by the Au neighbour atoms of the particle. These
results can shed light onto the question whether the particles are of “alloy” or
“core-shell” type. XPS data suggests an “alloy” (mixed) structure of the
particles.
9 Outlook
The aim of this thesis has been obtaining a better understanding of the
electronic and geometric properties of coinage (Ag and Au) metal nanoparticles
on HOPG and silica. Since only narrow size distributions of these particles
were prepared, depositing mass-selected clusters in a similar particle size
regime (i.e., 1-10 nm) could enable us to study changes of these properties with
more accurateness, i.e., atom by atom. With regard to smaller particle size
regimes (<1nm), a better geometric characterization by means of STM is still
necessary. Moreover, still remained is how studies on small mass-selected
clusters (< 15 atoms) obtained by our group somewhere else [178] can
complement the information acquired for larger particle size ranges in this
dissertation.
A possible subject of future studies can be the direct comparison of the
electronic structure of metal nanoparticles, with similar sizes but different
geometry. The chemical activity may be influenced by changes in the shape, for
example. Since the geometry of a supported nanoparticle is preferably to be
studied on a particle-by-particle basis, combined STM/STS experiments could
be appropriated for this purpose. In addition, the role of different substrates can
be further investigated. Especially, noble metal particles on different oxide
substrates are interesting, because of the usual enhance of their catalytic activity
due to these substrates.
Evidences of whether “alloy” or “core-shell” type formation of bimetallic
nanoparticles on different substrates take place (e.g. Ag-Cu or Au-Cu) can also
be further investigated. The interest on mixed-metal nanoparticles is based upon
the fact that catalytic activity in these systems is frequently found to be higher
than the corresponding monometallic catalysts [179]. Similar studies as those
reported in Chapter 7 about the size-dependent catalytic behaviors for the
oxidation and CO-oxidation are now in preparation for bimetallic (Ag-Au)
nanoparticles on HOPG in our group. A better characterization of the geometry
and growth mode of these complex bimetallic systems would still be necessary
for this purpose, especially in case of small particle size regimes (< 3 nm), in
which evidence of miscible bimetallic particles (alloy formation) has been
reported. A still unsolved problem is the control of the uniformity of the particle
in the preparation of miscible bimetallic particles [180], i.e., specific
117
118
9 Outlook
information about the composition (percent of mole fraction of atoms). To
achieve this, other synthetic routes in the preparation of these binary metallic
systems should be attempted [181, 182]
Most of the industrial applications in catalysis are at high pressure
environments (i.e. between 1 mbar and 1 bar). Usually the information about
the morphology, electronic structure and catalytic activity of these particles
under UHV conditions cannot be extrapolated to higher pressures. This problem
is known as “Pressure gap”. Thus, metal nanoparticles as model catalyst in
UHV conditions are often not applicable for the industrial catalysis. Techniques
such as TPD and STM in High-Pressure chamber can be applied to investigate
noble metal nanoparticles at higher pressure. In this context, High-Pressure
experiments of the oxidation and CO-oxidation of noble metal clusters are
already planned with a new Quadrupol Mass-Spectrometer in our laborory.
10 Zusammenfassung
Elektronische und geometrische Eigenschaften von Ag and Au
Nanopartikeln, die auf verschiedenen Substraten mittels Aufdampfen von
Metallatomen präpariert wurden, wurden bei Zimmertemperatur unter Ultra
Hoch Vakuum Bedingungen untersucht. Zu diesem Zweck wurden XPS (X-Ray
Photoelectron Spectroscopy) und STM (Scanning Tunneling Spectroscopy)
eingesetzt. Es wurden schmale Größenverteilungen der Partikel auf HOPG
(Highly Ordered Pyrolytic Graphite) und Silizium (SiO2/Si) Oberflächen
erreicht. Dies ermöglichte, die Zusammenhänge zwischen den Partikelgrößen
und den elektronischen und geometrischen Eigenschaften zu untersuchen. Eine
Charakterisierung dieser Eigenschaften in metallischen Nanopartikeln ist von
großer Bedeutung, da metallische Nanopartikel ideale Kandidaten für die
heterogene Katalyse verschiedener Reaktionen sind. Um die Frage zu
beantworten, wie die Natur des Substrates die unterschiedlichen Eigenschaften
der Partikel beeinflussen kann, wurden experimentelle Ergebnisse für die
Systeme Ag und Au Nanopartikel auf HOPG und SiO2/Si verglichen.
Zuerst wurde das System Ag/HOPG (Ag Nanopartikel auf HOPG) für
unterschiedliche Oberflächenbehandlungen erforscht. Nach Sputtern der HOPG
Oberfläche durch Ar-Ionen konnte eine viel höhere Ag Partikeldichte
beobachtet werden, was darauf hinweist, dass die Defektdichte eine wichtige
Rolle bei der Kernbildung der Partikel auf dem Substrat spielt. Die für dieses
System charakterisierten elektronischen und geometrischen Eigenschaften
wurden dann mit dem System Au/HOPG verglichen. Ein anderer Gegenstand
der Studie war das System Au/SiO2/Si. Die höhere Defektdichte der amorphen
Siliziumoxidschicht erhöht die Stabilität der auf dem Substrat gewachsenen
Partikel. Schließlich wurden bimetallische Ag-Au Partikel mit ähnlichen
Größen auf HOPG hergestellt und untersucht.
Bei allen Systemen wurden zunächst STM und anschließend XPS
Messungen durchgeführt. Die wichtigsten Ergebnisse für die untersuchte
System sind folgende:
Ag auf HOPG: Es wurde für Ag Nanopartikel auf HOPG ein
Wachstumsverhalten beobachtet, welches von anderen van der WaalsSubstraten (WSe2) abweicht. Das deutet darauf hin, dass zwischen den beiden
119
120
Zusammenfassung
Substraten subtile Unterschiede in der Substrat-Metall Wechselwirkung
vorhanden sind. Mit XPS wurden Änderungen in der elektronischen Struktur
der Ag Partikel mit abnehmender Teilchengröße gemessen (Verschiebung der
Energie des Ag 3d Kernniveaus). Die Energie shifts der Ag 3d Orbitale hängen
mit Anfangs- sowie Endzustandseffekten zusammen. Diese Ag 3d shifts
können sowohl durch die Teilchengröße als auch durch andere Faktoren wie
Partikelgeometrie und Substrat beeinflusst werden. Zusätzlich wurde eine
unterschiedliche Ag-Wolframoxid Wechselwirkung für Ag-Nanopartikel und Festkörper nachgewiesen.
Au auf HOPG: In diesem Fall wurden keine bedeutenden Unterschiede in
Partikelstruktur (Geometrie) und Wachstum von Au- und Ag-Partikeln auf
HOPG festgestellt. Es wurden dreidimensionales Wachstum und sechseckige
Formen mit vielen Stufen zwischen Terrassen beobachtet. Diese Information
weist darauf hin, dass das Partikelwachstum aufgrund der Zimmertemperatur
kinetisch beschränkt ist. Außerdem wurden unterschiedliche elektronische
Eigenschaften von Au und Ag Nanopartikeln auf HOPG gefunden. XPS
Untersuchungen von Au/HOPG bestätigten, dass Anfangszustandseffekte für
die Verschiebung der Energiezustände der Au 4f Orbitale verantwortlich sind.
Die wichtigsten Anfangszustandseffekte sind in diesem Fall die
Wechselwirkung zwischen Substrat und Partikel sowie der Einfluss von
unterkoordinierten Atomen. Der Vergleich von Ag und Au Nanopartikeln auf
HOPG mittels einer Auger Analyse weist darauf hin, dass die Ag und Au
Nanopartikel aufgrund eines Ladungstransfers zwischen dem Substrat und den
Partikeln jeweils positiv oder negativ geladen sind.
Au auf SiO2/Si: Hier wurde beobachtet, dass Au Nanopartikel auf SiO2/Si
dreidimensionales Wachstum zeigen. Mittels XPS wurde festgestellt, dass die
elektronischen Eigenschaften von Au Nanopartikeln (d.h. kleiner als ~10
Nanometer) vom Au Festkörper abweichen. Verglichen mit Au Nanopartikeln
auf HOPG, zeigen die Partikel auf SiO2/Si höhere Core Level Shifts des Au 4f
Zustandes, die ähnlich zu Core Level Shifts sind, die für schlecht leitende
Substrate (z.B Aluminium- und Titanoxide) gefunden wurden. Außerdem
wurde gezeigt, dass eine Abnahme der durchschnittlichen Koordinationszahl
der Atome innerhalb der Au Nanopartikel auf SiO2/Si die elektronische Struktur
des Valenzbandes beeinflusst.
Bimetallische Ag-Au Nanopartikel auf HOPG: Es wurde gezeigt, dass
bimetallische Ag-Au Nanopartikel auf HOPG mit relativ schmalen
Größenverteilungen präpariert werden können. XPS Studien deuten darauf hin,
Zusammenfassung
121
dass die Ag 3d Zustände in diesen Partikeln durch die benachbarten Au Atome
innerhalb des Partikels beeinflusst werden können. Diese Untersuchungen
dienen dazu herauszufinden, ob die bimetallischen Partikel als „Legierung“
oder als „Core-Shell“ vorliegen. Die XPS Daten weisen auf eine „Legierung“
(Gemisch) hin.
122
A List of Figures
Fig 2.1
Catalytic activity and electronic structure of gold nanoparticles on
TiO2 (100) ...........................................................................................6
Fig 2.2
Mass spectra of free Au-n cluster anions............................................7
Fig 2.3
XPS spectra of Aun cluster (n=2 - 10) deposited on SiO2. .................8
Fig 2.4
Percent of edge Au atoms by particle as a function of particle size .11
Fig 2.5
Interaction energy per molecule for CO and O adsorption versus
coordination number for gold atoms in various geometries.............11
Fig 2.6
Two STM images of Au+3 clusters on TiO2 (100)-(1x1) surface at
room temperature..............................................................................12
Fig 3.1
Main principle of STM. A feedback mechanism for the piezo controls
the distance between tip and sample. Tunneling current and voltage
can be monitored as the tip scans over the surface. .........................17
Fig 3.2
Rectangular barrier with a potential height Φ.. ...............................18
Fig 3.3
Energy band diagram of a STM tunnel junction (a) before applying a
bias voltage at equilibrium and (b) after applying a positive voltage
relative to the sample. In this case, an electron current flows from tip
to sample. ..........................................................................................19
Fig 3.4
Schematic energy diagram of electron tunneling with respect to the
density of states of the sample. The occupied states of the sample
(indicated with dark color) generate the current []..........................21
Fig 4.1
"Universal curve" of the electron Inelastic Mean Free Path
(IMFP=λ) versus kinetic energies for different materials ................23
Fig 4.2
Schematic energy diagram of X-ray photoelectron spectra..............24
Fig 4.3
XPS spectrum obtained from a Pd bulk sample . ..............................27
Fig 4.4
Energetic scheme of the Auger process KL1L2,3 carried out using X-
123
124
A List of Figures
ray radiation ..................................................................................... 28
Fig 4.5
XPS spectrum of the C1s state of a HOPG surface. ......................... 30
Fig 5.1
Sample-holder used in the UHV preparation chamber .................... 32
Fig 5.2
Two images of a Pt/Ir tip .................................................................. 32
Fig 5.3
Scheme of the experimental procedure............................................. 34
Fig 5.4
Three dimensional drawing of the UHV system (2).......................... 35
Fig 5.5
Picture of the STM stage taken from top view. Main components of
the Omicron STM1 are indicated. .................................................... 36
Fig 5.6
Left: Schematic draw of a STM Tip .................................................. 37
Fig 5.7
Schematic diagram of the experimental equipment in the UHV XPS
chamber ............................................................................................ 38
Fig 5.8
Schematic diagram of a Concentric Hemispherical Analyzer (CHA)
for the detection of the photoelectron energy in XPS ....................... 39
Fig 5.9
(a) Schematic diagram of the lattice structure of a small piece of
HOPG with 3 atomic-layers. (b) STM image with atomic resolution of
HOPG (2.2 x 1.1 nm2)....................................................................... 41
Fig 5.10 HREELS Spectra of a non-sputtered HOPG surface ....................... 42
Fig 5.11 Two different STM images of the same sample of Ag nanoparticles
deposited on non-sputtered HOPG (500 x 500 nm2), ...................... 42
Fig 5.12 HREELS spectrum of a SiO2 surface before annealing ................... 43
Fig 6.1
Two STM images of the same sample for Ag nanoparticles on HOPG;
Upper left: 135.3 x 132.3 nm2. Upper right: 50 x 50 nm2. ............... 46
Fig 6.2
Two STM images of the same sample for Ag nanoparticles on HOPG;
upper Left: 600x 320 nm2. Upper right: 214.5 x 214.5 nm2 ............. 47
Fig 6.3
Two STM images of the same sample of Ag nanoparticles on
sputtered HOPG; upper left: 352.3 x 352.3 nm2. Upper right: 154.9 x
154.9 nm2 .......................................................................................... 48
A List of Figures
125
Fig 6.4
Two STM images of Ag nanoparticles on sputtered HOPG; Upper
left: 394.4 x 349.8 nm2. Upper right: 45 x 45 nm2. ...........................49
Fig 6.5
Particle diameter versus particle height of Ag nanoparticles grown
on HOPG...........................................................................................50
Fig 6.6
STM image of a sample for Au nanoparticles on HOPG:(300 x 300
nm2) ...................................................................................................51
Fig 6.7
Two STM images of the same sample for Au nanoparticles on
HOPG; Upper left: 3-D image (256.5 x 256.5 nm2). Upper right: 2-D
image (256.5 x 256.5 nm2). ...............................................................52
Fig 6.8
Two STM image of the same sample for Au nanoparticles on HOPG;
Upper left: 430 x 430 nm2; Upper right: 95 x 145 nm2 ....................53
Fig 6.9
STM image of Au nanoparticles (nanostructured film) grown on
HOPG (350 x 500 nm2). Tunneling parameters: -1.2 V, 1.3 nA. ......54
Fig 6.10 Upper picture:STM image of Au nanoparticles grown on HOPG (220
x 220 nm2 Lower picture: Step profiles.............................................55
Fig 6.11 Au particle diameter versus particles height. ...................................56
Fig 6.12 Left: STM image of Ag nanoparticles grown on HOPG (50 x 50 nm2)
Tunneling parameters: 3.0 V, 0.7 nA. Right: FWHM of the particle
profile is shown in the plot. ...............................................................57
Fig 6.13 Two STM images of the same sample for Ag nanoparticles on nonsputtered HOPG. The images were taken at different spots of the
sample. Left: 500 x 500 nm2. Right: 495 x 335 nm2..........................60
Fig 6.14 Time successive STM images at same location for a HOPG sample on
which Ag nanoparticles were grown. Tunnelling parameters: -3 V,
0,7 nA. a): 150 x 115 nm2; b): 146 x 114 nm2 (6 minutes after a); c):
147 x 109 nm2 (8 minutes after b). ....................................................61
Fig 6.15 Left: STM image of Ag nanoparticles on sputtered HOPG (214.5 x
214.5 nm2); Right: Ag nanoparticles on non-sputtered HOPG with
same Ag exposure time (30 min); (207.2 x 181.4 nm2) .....................63
Fig 6.16 Two STM images of the same sample of Ag nanoparticles on a non-
126
A List of Figures
sputtered HOPG surface, demonstrating that Ag preferentially
decorates step edges. Left: 871.1 x 871.1nm2, Right: 246 x 246 nm2.
Tunneling parameters: 0.2 V, 1 nA................................................... 63
Fig 6.17 Schematic diagram of the three growth modes usually observed for
metal overlayers [80]. ...................................................................... 65
Fig 6.18 Three dimensional STM image of Ag nanoparticles on HOPG with a
particle size distribution of 3 nm in diameter and 1 nm in height. (50
x 50 nm2). Tunneling parameters: 3.0 V, 0.7 nA. ............................. 65
Fig 6.19 Atomic surfaces processes and characteristic energies, which
participate in nucleation and film growth. ....................................... 65
Fig 6.20 Left: One dimensional representation of the diffusion barrier as a
function of position on an ideal defect-free surface. Right:
Geometrical corrugation shown by a three dimensional STM image
with atomic resolution of a defect-free HOPG surface (1.7 x 1.0 nm2;
0.1 V, 0.7 nA). ................................................................................... 66
Fig 6.21 a) STM image of 2 Au nanostructured particles (60 x 60 nm2). b)
Schematic draw of the facet structure of the particle; Hexagonal
terraces with plane (111) connected with step planes (100) and (111)
form the particles. c) Schematic draw of atoms diffusing on the
surface subject to an Ehrlich-Schwöbel barrier. .............................. 67
Fig 6.22 Left: STM image of a single Au nanoparticle (44 x 44 nm2).
Tunnelling parameters:1.5 V, 0.2 nA. Right: Step profile of the blue
line. Atomic steps of ~0.2 nm size can be recognized....................... 67
Fig 6.23 Illustration of a kink of a nanoparticle formed by facets (100) and
(111). Above these particles a STM tip with two apexs scans the
particle from right to the left............................................................ 69
Fig 6.24 a) XPS Ag 3d core level of Ag nanoparticles on HOPG as a function
of the particle size. b) Summary of Ag 3d level shifts and broadening
of Ag 3d peaks as a function of the particle size............................... 72
Fig 6.25 Schematic energy level diagram of metal atoms, dimers, clusters
(nanoparticles) and bulk. As the cluster size increases, a continuous
metallic band develops].................................................................... 74
A List of Figures
127
Fig 6.26 XPS Au 4f levels as function of particle size a) For the samples
already imaged by STM. b) For the particle size regime from bulk to
less than ~ 1.5 nm. The botton 5 Au 4f spectras are presented again in
detail in Fig 6.27. ..............................................................................76
Fig 6.27 XPS spectra of Au 4f core levels as a function of particle size. An
inversion in the core level shift can be observed in the smallest Au
particle size regime (<1.5 nm). .........................................................77
Fig 6.28 Top: XPS Ag 3d core level as a function of the particle size. Botton:
XPS Ag 4d valence levels as function of the particle size. For each
spectrum, different data acquisition times were used, which are
denoted. Note: Low intensities of Ag 4d peaks in comparison with Ag
3d were observed. .............................................................................82
Fig 6.29 Ag 3d and Ag 4d states as a function of Ag coverage (i.e. particle
size) on HOPG. The changes of the centers of 3d and 4d states
confirm a deviation of only ~ 0.2 eV. ................................................83
Fig 6.30 Ag MVV Auger peaks as a function of particle size. .........................84
Fig 6.31 Final and initial state contribution as a function of Ag 3d core level
shifts (i.e particle size). .....................................................................85
Fig 6.32 Changes of the Ag 3d core levels with increasing W-oxide coating
thickness for Ag nanoparticles on HOPG with a mean diameter of 4
nm......................................................................................................88
Fig 6.33 Changes of the Ag 3d core levels with increasing W-oxide coating
thickness for Ag bulk. ........................................................................89
Fig 6.34 W core levels showing formation of W- oxide coating. The W 4f levels
are shifted to higher binding energies compared to those of metallic
W, indicating W-oxide formation. .....................................................90
Fig 6.35 STM image of Au nanoparticles on native silica layers on Si wafer
(193 x 193 nm2). Tunneling parameters: 2V, 0.1 nA. Lower pictures:
Diameter and height distribution of the particles.............................91
Fig 6.36 STM image of Au nanoparticles on native silica layers on Si wafer
(124 x 124 nm2). Tunneling parameters: 4.7 V, 0.1 nA. Lower
pictures: Diameter and height distribution of the particles..............92
128
A List of Figures
Fig 6.37 STM image of Au nanoparticles on native silica layers on Si wafers
(123 x123 nm2). Tunnelling parameters: 4.7 V, 0.1 nA. Lower
pictures: Diameter and height distribution of the particles. ............ 93
Fig 6. 38 Particle diameter versus particle height of Ag nanoparticles grown
on Silica ............................................................................................ 94
Fig 6.39 XPS Au 4f levels for Au nanoparticles on SiO2/Si. From a to h, Au
coverage increases. A Si satellite peak with decreasing particle size
appear. For each spectrum, different data acquisition times were
used, which are denoted. .................................................................. 95
Fig 6.40 Centers of Au 4f7/2 peaks as a function of amount of Au deposite (i.e.
particle size)...................................................................................... 96
Fig 6.41 a) XPS spectra of Si 2p level as a function of Au coverage. b) Centers
of the Si 2p peaks as a function of amount of Au deposited.............. 97
Fig 6.42 Valence band structures as a function of Au coverage. A spectrum of
a bare Si wafer was subtracted from each spectrum. The five spectra
were taken from the samples of the spectra h-d in Fig 6.39. One can
observe the band narrowing and the decrease of the splitting between
the peaks A and B. ............................................................................ 99
Fig 6.43 STM image of bimetallic nanoparticles (Ag@Au) grown on HOPG
(223.6 x 223.6 nm2). Tunnelling parameters: 1.78 V, 0.15 nA. Lower
picture: Size distribution of the particles........................................ 101
Fig 6.44 STM image of bimetallic nanoparticles (Ag@Au) grown on HOPG
(194.2 x 194.2 nm2). Tunnelling parameters: -1.89 V, 0.045 nA.
Lower picture: Diameter and height distribution of particles........ 102
Fig 6.45 Two STM images of bimetallic nanoparticles (Au@Ag) on HOPG.
Left: 518 x 518 nm2. Right: 225 x 225 nm2. Tunnelling parameters: 1.97 V, 0.08 nA. Lower picture: Diameter and height distribution of
the particles. ................................................................................... 103
Fig 6.46 Two STM images of the same sample of bimetallic nanoparticles
(Ag@Au) grown on non-sputtered HOPG. Left) 500 x 500 nm2; STM
parameters: -3.7 V, 0.88 nA. Right) 233.6 x 233.6 nm2; STM
parameters: -3.7 V, 0.33 nA. Note: Especial structures are formed,
A List of Figures
129
compound of several single particles..............................................104
Fig 6.47 STM image of bimetallic nanoparticles (Ag@Au) grown on slightly
sputtered HOPG: 491.2 x 320 nm2. Tunneling parameters: - 3.7 V,
0.8 nA. Lower picture: Diameter and height distribution of the
particles. Broader particle size distribution was found compared with
samples of Fig 6.43-6.45.................................................................105
Fig 6. 48 Particle diameter versus particle height of Ag nanoparticles grown
on Silica...........................................................................................105
Fig 6.49 a) XPS spectra of Ag 3d states of (Ag-Au) nanoparticles as a function
of particle size and. b) Summary of the Ag 3d level shifts of Ag
nanoparticles and bimetallic (Ag-Au) particles as a function of
particle size. ....................................................................................108
Fig 7.1
XPS spectra of O 1s states for Ag nanoparticles on HOPG (~ 6 nm)
after exposition to atomic oxygen and subsequent reduction by CO at
room temperature............................................................................111
Fig 7.2
XPS spectra of Ag 3d level for Ag nanoparticle (~6 nm) after
exposing to atomic oxygen and subsequent reduction by CO.........112
Fig 7.3
XPS spectra of O 1s states of Au nanoparticles with a mean size of 5
nm after expostion to atomic oxygen and subsequent reduction by CO
.........................................................................................................113
Fig 7.4
Au 4f states for Au nanoparticles with a mean size of 5 nm after
exposition to atomic oxygen and subsequent reduction to CO .......113
Fig 7.5
XPS spectra of Au 4f levels of Au nanoparticles (between ~1.4 and ~3
nm size) grown on SiO2/Si. Au nanoparticles were exposed to atomic
oxygen and subsequently reduced by CO........................................114
130
B List of Presentations
Oral presentation in Conferences:
1
Plasmonenanregungen in Ag-Clustern verschiedener Größe auf Graphit:
I. Lopez-Salido, D. Stolcic, N. Beltram, Y. D. Kim, G. Ganteför,
Spring Meeting of the German Physical Society, Dresden, March, 24-28
(2003).
2
Electronic and Geometric Properties of Au and Ag Nanoparticles:
I. Lopez-Salido, D. C. Lim, R. Dietsche, M. Bubek, Y. D. Kim, G.
Ganteför,
Spring Meeting of the German Physical Society, Dresden March, 26-31
(2006).
3
Electronic and Chemical Properties of Ag and Au Nanoparticles:
I. Lopez-Salido, D. C. Lim, R. Dietsche, M. Bubek, Y. D. Kim, G.
Ganteför,
SFB 513-Klausurtagung in Söllerhaus, Austria, 18-20 September (2006).
Posters presented in Conferences:
•
Plasmon Investigation of metal Nanoparticles on HOPG with HREELS:
I. Lopez-Salido, N. Bertram, D. Stolcic, Y. D. Kim, G. Ganteför:
513-Workshop& Krupp Symposium, Konstanz, 6-9. July (2003)
•
Anomalous Behaviour of small metal Clusters on HOPG, Studies on
Plasmon Resonances of Ag Clusters using HREELS:
I. Lopez-Salido, N. Bertram, D. Stolcic, Y. D. Kim, G. Ganteför
Wilhelm und Else Heraeus Seminar: “Freie und deponierte Cluster” 5-10.
October (2003), Brand (Österreich).
•
Plasmonenanregungen in Ag-Clustern verschiedener Größe auf Graphit:
I. Lopez-Salido, N. Bertram, Y. D. Kim and G. Ganteför.
Spring Meeting of the German Physical Society, Dresden, March, 24-28
131
132
B List of Presentations
(2003).
•
Electronic Structure and Chemical Properties of metal Clusters in various
environments:
D.C. Lim , I. Lopez-Salido, N. Bertram, Felix v.Gynz-Rekowski, , Tim
Fischer R. Dietsche, Y. D. Kim, G. Ganteför.
Sonderforschungsbereich (SFB) 513, Nanostrukturen an Grenzflächen und
Oberflächen Begutachtung 2004, 5-8 October (2004), Konstanz.
•
Electronic, Geometric and Chemical Properties of Au and Ag metal
Nanoparticles and Clusters in various environments:
I. Lopez-Salido, D. C. Lim, R. Dietsche, M. Bubek, G. Ganteför and Y. D.
Kim
Deutsche Cluster Treffen, Bad Herrenalb, 11-16. September (2005).
•
Oxidation and Reduction of metal Nanoparticles:
D. C. Lim , I. Lopez-Salido, , R. Dietsche, M. Bubek, G. Ganteför and Y. D.
Kim
Trends in Nanoscience, Kloster Irsee, 8-12, October (2005).
•
Interaction with Oxygen of Ag mass-selected Clusters and Nanoparticles:
I. Lopez-Salido, D. C. Lim, T. Fischer, F. V. Gynz-Rekowski, G. Ganteför,
Y. D. Kim
S3C 2005- Symposium on Size-Selected clusters, Brand (Austria), from 28th
February until 3th May (2005).
•
Growth and Properties of Ag Nanoparticles on HOPG:
I. Lopez-Salido, D. C. Lim, T.Fischer, R. Dietsche, and Young Dok Kim.
Spring Meeting of the German Physical Society, Berlin, March, 24-28
(2005).
•
Oxidation and Reduction of metal Nanoparticles:
Dong Chan Lim, Ignacio Lopez-Salido, , Rainer Dietsche, Moritz Bubek, G.
Ganteför and Young Dok Kim
Spring Meeting of the German Physical Society, Dresden, March, 24-30
(2006).
C References
[1]
D. M-K. Lam, B. W. Rossiter, Sci. Am. 265, 80 (1991).
[2]
Leung C, Xirouchaki C, Berovic N and Palmer, Adv. Mater. 16, 223,
(2004).
[3]
Liang Z J, Susha A and Caruso F. Chem. Mater.15 3176, (2003).
[4]
C. B. Murray, S. Sun, H. Doyle, T. Betley, Mater. Res. Soc. Bull. 26,
985 (2001).
[5]
P. V. Kamat et.al., J. Phys. Chem. B 106, 7729 (2002).
[6]
Haruta. M, Date. M, Applied Catalysis A. General, Vol 222, 1 (2001),
pp. 427-437.
[7]
Jacob I Kleiman, Protection of Space Materials from the Space
Environment. Editor Springer, (2001)
[8]
C. T. Campbell and M. T. Paffett, Surf. Sci. 139, 396. (1984).
[9]
B. Grant and R. M. Lambert, J. Catal. 92, 364, (1985).
[10]
Sondi I, Salopek-Sondi B, J. Colloid Interface Sci. (2004); 275(1):177182.
[11]
Cancer Research, 65 (12), June 15, 2005
[12]
Christof M. Niemeyer, Angewandte Chemie International, Vol 40, Issue
22, 4128-4158 (2001).
[13]
Yugang Sun, Younan Xia, Science 13, Vol. 298, pp, 2176-2179 (2002).
[14]
M. Haruta, Catalysis Today 36 (1997) 153.
[15]
M. Valden, X. Lai, D.W. Goodman, Science 281 (1998), 1647.
[16]
M. Haruta, S. Tsubota, T. Kobayashi, H. Kageyama, M.J. Genet, B.
Delmon; J. Catal. 144 (1993), 175.
[17]
N. Lopez, J.K. Norskov, J. Am. Chem. Soc. 124 (2002), 11262.
[18]
A. Sanchez, S. Abbet, U. Heiz, W.-D. Schneider, H. Häkkinen, R.N.
Barnett, U. Landman, J. Phys. Chem. A 103 (1999), 9573.
[19]
U. Heiz, F. Vanolli, L. Trento, and W.-D. Schneider, Rev. Sci. Instrum.
68, (1986-1997).
133
134
C References
[20]
T. Schlenker, Diploma Thesis, Universität Konstanz, 2000.
[21]
H. Haberland, Clusters of Atoms and Molecules I & II, Springer, Berlin,
(1995).
[22]
Felix Von Gynz-Rekowski, Sauerstoffadsorption an freien und
deponierten Clustern. Dissertation, Universität Konstanz (2005).
[23]
M. C. Bartel and J. W. Evans, Phys. Rev. B 46 (1992)
[24]
V.M. Garamus, T. Maksimova,W. Richtering, C. Aymonier, R.
Thomann,L. Antonietti, S. Mecking, Macromolecules 37 (2004), 7893.
[25]
D. C. Lim, Y. D. Kim, I. Lopez-Salido, accepted in Apl. Surf. Science
(2006).
[26]
Charles T. Campbell, Science (2004), Vol. 306, 5694, 234-235.
[27]
S. B. Dicenzo et.al., Phy. Rev. B, (1988), 289-329.
[28]
M. M. Alvarez, J. T. Khoury, T. G. Schaaff, M. N. Shafigullin, I,
Vezmar and R. L. Whetten, J. Phys. Chem. B. (1997), 101, 3706.
[29]
NT.V. W. Janssens, B. S. Clausen, Y. Xu, M. Mavrikakis, T. Bligaard,
J. K. Norskov, Jounal of Catalysis 223 (2004), 232.
[30]
Häkkinen, H.; Moseler, M.; Landman, U; Phys. Rev. Lett (2002), 89,
033401.
[31]
W. F. Egelhoff. Jr, et. al.,Surf. Sci. Rep. 6, 253 (1987).
[32]
G. K.Wertheim, Z. Phys. B 66, 53 (1987); Z. Phys. D 12, 319 (1989).
[33]
G. K. Wertheim, S. B. DiCenzo, and D. N. E. Buchanan, Phys. Rev. B
33, 5384 (1986).
[34]
F. Parmigiani et al., J. Electron. Spectrosc. Relat. Phenom. 50, 39
(1990).
[35]
M.G. Mason, Phys. Rev. B 27, 748 (1983).
[36]
A. Sandell et al., J. Vac. Sci. Technol. A 14, 1546 (1996).
[37]
P. S. Bagus et al., J. Electron Spectrosc. Relat. Phenom.100, 215 (1999).
[38]
G. C. Bond, D.T. Thompson, Gold Bull. 33 (2000) 41.
[39]
R. E. Watson and M. L. Perlman, Struc. Bonding (Berlin) 24, 83 (1975).
[40]
Haberland. H, Moseler. M, Qiang. Y, Rattunde. O, Reiners. T and
Thurner. Y, Phy. Rev. Lett; 3, 887 (1996).
C References
135
[41]
G. Ganteför, HR. Siekmann, HO. Lutz, KH. Meiwes-Broer; Chemical
Physics Letters, 165, Issue 4, p.293-296 (1990).
[42]
Richard E. Smalley et al., Acc. Chem. Res., Vol. 25, No. 3, (1992)
[43]
Jun Li, Xi Li, Hua-Jin Zhai, Lai-Sheng Wang, Science, Vol 299, (2003).
[44]
H.-G. Boyen, G. Kästle, F. Weigl, and P. Ziemann, G. Schmid, M.G.
Garnier, P. Oelhafen; Phy. Rev. Lett. 87 (2001), 276401.
[45]
H. Boyen, G. Kästle, F. Weigl, G. Schmid, M. Ganrier, P. Oelhafen, et
al., Science, 297:1533–1536, (2002).
[46]
D. Stolcic, M. Fischer, Y. Kim, Q. Sun, P. Jena, und G. Ganteför, J. Am.
Chem. Soc. 125: 2848–2849, (2003).
[47]
B. E. Salisbury, W. T. Wallace, R. L. Whetten, Chem. Phys. 262, 131
(2000).
[48]
L. Wöste et al., Phys. Chem. Chem. Phys. 7, 2706 (2005).
[49]
M. Niemietz, P. Gerhardt, G. Ganteför and Y. D. Kim, Chem. Phys.
Lett. 380, 99 (2003).
[50]
M. Niemietz, M. Engelke, G. Ganteför; Photodesorption of O2 from
Ag2-: A Time-Resolved Study of Ag2O2-. Unpublished Results (2006).
[51]
Moritz Bubek. Diplomarbeit. Photoelektron Spectroscopie an Gold
Clustern, University Konstanz (2006).
[52]
D. C. Lim, M. Bubek, R. Dietsche, G. Ganteför,Y. D. Kim, Chem. Phys.
Chem. Vol 7, Issue 9, (2006), 1909-1911.
[53]
Sungsik Lee, Chaoyang Fan, Tanpin Wu, and Scott L Anderson, J. Am,
Chem. Soc. (2004), 126, 5682-5683.
[54]
Y. D. Kim, D. C. Lim, I. L .Salido, R. Dietsche, M. Bubek, Angew.
Chem. Int. Ed., 45, 1-4 (2006).
[55]
V.A. Bondzie, S.C. Parker, and C. T. Campbell, J. Vac. Sci. Techn. A,
(1999), 17, 1717.
[56]
U. Heiz et al., Science 307, 403 (2005).
[57]
A. L. de Oliveira, A. Wolf and F. Schüth, Catalysis Letters. 73 (2-4), 57
(2001).
[58]
E. Wahlström, N. Lopez, J. K. Norkov, F. Besenbacher et al., Phy. Rev.
Lett, Vol 90, 2 (2003).
136
C References
[59]
J. D. Grünwaldt, C. Kiener, C. Wögerbauer, and A. Baiker, J. Catal.,
(1999), 181, 223.
[60]
H.Häkkinen, S. Abbet, A. Sanchez, U. Heiz and U. Landman; Angew.
Chem., Int. Ed., (2003), 42, 1297–1300.
[61]
M. Haruta, CATTECH 6, 102 (2002).
[62]
M. Mavrikakis, P. Stoltze and J. Nørskov, Catal. Lett., (2001), 64, 101.
[63]
R. Meyer, C. Lemire, Sh. K. Shaikhutdinov, H.-J. Freund; Surface
Chemistry of Catalysis in Gold. Gold. Boll. (2004).
[64]
S. R. Bahn, N. Lopez, J.K. Norskov and K. W. Jacobsen, Phy. Rev. B,
(2002), 66, 081405.
[65]
Lauren Benz, Xiao Tong, Michael T. Bowers, and Steven K. Buratto et.
al., J. Chem. Phys. 122, 081102 (2005).
[66]
M T. Bowers, K. Buratto et al., J. Am. Chem. Soc. (2005), 127, 1351613518.
[67]
Karsten Bromann, Harald Brune, Klaus Kern et.al., Surf Science 377379 (1997) 1051-1055.
[68]
R. Kligeler, P. S. Bechthold, M. Neeb, and W. Eberhardt, Review of
Scientific instruments, (2002), Vol 73, number 4.
[69]
U. Heiz, W-D Schneider. Critical Reviews in Solid State and Materials
Sciences, 26 (4):251–290 (2001).
[70]
F. Claeyssens, S Pratontep, C Xirouchaki and R.E. Palmer,
Nanotecnology, 17 (2006) 805-807.
[71]
The Chemical Record, Vol 2, Issue 6 , Pages 446- 457 (2002).
[72]
N. Nilius, N. Ernst, and H–J. Freund, Phys Rev Lett, Vol 84, 17 (2000).
[73]
H. Hövel et al., Appl. Phys. A. 72, 295–302 (2001).
[74]
M. Amman, R. Wilkins, E. Ben-Jakob, P.D. Maker, R.C. Jaklevic; Phys.
Rev. B 43, 1146 (1991).
[75]
R.P. Andres, T. Bein, M. Dorogi, S. Feng, J.I. Henderson, C.P.
Kubiak,W. Mahoney, R.G. Osifchin, R. Reifenberger: Science 272,
1323 (1996).
[76]
W. Schneider et al., Phys Rev Lett, Vol 80, n 13 (1998).
C References
137
[77]
G. Binning, H. Rohrer, Ch. Gerber, E. Weibel: Surface Studies by
Scanning Tunneling Microscopy. Phys. Rev. Lett. 49, 57-61 (1982).
[78]
G. Binnig, H. Rohrer, C. Gerber, and E. Weibel. Tunneling through a
controllable vacuum gap. Applied Physics Letters, 40:178_180, (1982).
[79]
M. C. Desjonquéres and D. Spanjaard, Concepts in Surface Physics, 2nd
Edition, Springer, Berlin, (1996).
[80]
J. Venerables, Introduction to Surface and Thin Film processes,
Cambridge University press, (2001).
[81]
R. Wiesendanger, "Scanning Probe Microscopy and Spectroscopy:
Methods and Applications", Cambridge University Press, (1998).
[82]
D. Bonnell, "Scanning Probe Microscopy and Spectroscopy: Theory,
Techniques, and Applications", 2nd ed, Wiley-VCH, New York, (2001).
[83]
H.-J. Guentherodt, R. Wiesendanger, "Scanning Tunneling
Microscopy", Vol. I, II, and III, Springer, (1993, 1995, 1996).
[84]
http://www.almaden.ibm.com/vis/stm/
[85]
John G. Simmons, J. of App. Phys. 34, 1793 (1963).
[86]
D. J. Griffiths, Introduction to Quantum Mechanics, Prentice-Hall,
Englewood Cliffs, NJ (1995).
[87]
N Zettili, “Quantum Mechanics Concepts and Applications”, John
Wiley and Sons (2001).
[88]
G. Binnig and H. Rohrer, Helv. Phys. Acta. 55, 726 (1982).
[89]
G. Binnig and Rohrer, Surf. Sci. 152-153, 17 (1985).
[90]
J. Tersoff and D.R. Hamann, Phys. Rev. Lett. 50, 1998 (1983).
[91]
J. Tersoff and D.R. Hamann, Phys. Rev. B, 31, 805 (1985).
[92]
http://www.fkp.uni-erlangen.de/methoden/stmtutor/stmpage.html
[93]
Schneider, W.-D., (ed.), Scanning Tunneling Spectroscopy, J. El. Spec.,
109, (2000).
[94]
M. Bode, S. Heinze, A. Kubetzka, O. Pietzsch, X. Nie, G. Bihlmayer,
and R. Wiesendanger, “Magnetization-Direction-Dependent Local
Electronic Structure Probed by Scanning Tunneling Spectroscopy”,.
Phys. Rev. Lett., 89 (23): (2002).
138
C References
[95]
O. Pietzsch, A. Kubetzka, M. Bode, and R. Wiesendanger; “Observation
of Magnetic Hysteresis at the Nanometer Scale by Spin-Polarized
Scanning Tunneling Spectroscopy”; Science, 292: 2053-2056, (2001).
[96]
Ali Yazdani, B.A. Jones, C.P. Lutz, M.F. Crommie, and D.M. Eigler,
“Probing the Local Effects of Magnetic Impurities on
Superconductivity”, Science, 275: 1767-1770, (1999).
[97]
Anatoli Polkovnikov, Subir Sachdev, and Matthias Vojta; “Impurity in a
d-Wave Superconductor: Kondo Effect and STM Spectra”, Phys. Rev.
Lett., 86(2): 296, (2001).
[98]
A. Einstein, „Uber einen die Erzeugung und Verwandlung des Lichtes
betreffenden heuristischen Gesichtspunkt. Annalen der Physik“, 7:132–
148, 819059 (1905).
[99]
T.E. Madey and J.T.Yates, J. Vac. Sci. Techn. 8 (1971) 525.
[100] G. Ertl, K. Küppers, Low Energy Electrons and Surface Chemistry,
VCH Verlagsgesellschaft, Weinheim, (1985).
[101] http://www.bessy.de.
[102] Koopmans, T. "Ordering of Wave Functions and Eigenvalues to the
Individual Electrons of an Atom.", Physica (1933), 1, 104.
[103] D. C. Lim, I. Lopez-Salido and Y. D. Kim, Surf. Sci., 2005, Vol. 598,
Issues 1-3, 96-103.
[104] D. C. Lim, I. L.-Salido and Y. D. Kim, Surf. Sci. 600 (2006) 507-513.
[105] Bagus, P. S.; Brundle, C. R.; Pacchioni, G.; Parmegiani, F. Surf. Sci.
Rep. 1993, 19, 265.
[106] K. Luo, X. Lai, C.-W. YA. Davis, K.K. Gath, and D. W. Goodman, J.
Phys. Chem. B, Vol. 109, No. 9, (2005).
[107] Tougard and Sigmund: Phys. Rev. B 25, 4452 (1982).
[108] Handbook of the Elements and Native Oxide, XPS International, Inc
(1999).
[109] http://www.chem.qmul.ac.uk/surfaces/scc/
[110] D. Briggs und M. P. Seah, “Practical Surface Analysis Volume I“, John
Wiley & Sons, 2. Ed. (1990).
[111] B. Klipp. Deposition massenselektierter Aluminiumcluster. Dissertation,
Universität Konstanz (2000).
C References
139
[112] H. Ballot, Personal Staff, Group Prof. Leiderer from University
Konstanz (2005).
[113] Hövel, H.; Becker, Th.; Bettac, A.; Reihl, B.; Tschudy, M.;Williams, E.
J. J. Appl. Phys. (1997), 81, 154-158.
[114] H.-J. Freund; F. Besenbacher, et al. , Phys. Rev. Lett. 83, 4120 (1999).
[115] H. Ibach, ”Topics in Current Physics 4: Electron spectroscopy for
Surface Analysis”. Springer Verlag, Berlin, Heidelberg, New York
(1997).
[116] Omicron. EA 125 Energy Analyzer, Technical Reference Manual,
(1997).
[117] I. Lopez-Salido, D. C; Lim, R. Dietsche, Y. D. Kim, J. Phys. Chem. B,
110, 128-10136 (2006).
[118] Bukhtiyarov, V. I.; Carley, A. F.; Dollard, L. A.; Roberts, M. W.Surf.
Sci. (1997), 381, L605-L608.
[119] Bukhtiyarov, V. I.; Kaichev, V., J. Mol. Catal. A: Chem. (2000), 158,
167-172.
[120] D. Stolcić, HREELS; unpublished results from AG. Gantefoer,
Universität Konstanz. (2002).
[121] A. Rettenberger, P. Bruker, M. Metzler, F. Mugele, Th. W. Matthes, M.
Böhmisch, J. Boneberg, K. Friemelt, P. Leiderer, Surf. Sci. 402–404
(1998) 409; J. Zimmermann, Dissertation, University of Konstanz,
(2004).
[122] T. Irawani. Barkeh. Hövel et al, Appl. Phys. A 80, 929–935 (2005).
[123] Zeitschrift für Kristallographie (Bauer 1958)
[124] Overbury, S. H.; Bertrand, P. A. ; Somorjai, G. A. Chem. Rev. (1975),
75, 547
[125] V. E. Heirich and P. A. Cox; The Surface Science of Metal Oxide,
Cambridge University Press, (2000).
[126] Mao, C. –F. ; Vannice, M . Acata., J. Catal. (1991), 71, 247.
[127] Mezey, L. z.; Giber, J. Jpn. J. Appl. Phys. (1982), 21, 1569.
[128] H. P. Bonzel: Phys. Rep. 385, 1 (2003).
[129] F. Besenbacher, J.-Freund et al. Phys Rev Lett, vol 83, 20 (1999).
140
C References
[130] Phys. Rev. B Condens. Matter (1996), 15; 54(24):17930-17937.
[131] Kinetics of nucleation in surfaces mediated epitaxy. Phys. Rev. B
Condens. Matter. (1996); 53(7): 4148-4155.
[132] Caspersen, K. J.; Stoldt, C. R.; Layson, A. R.; Bartelt, M. C.; Thiel, P.
A.; Evans, J. W., Phys. Rev. B (2001), 63, 085401.
[133] K. Luo, T.P.St. Clair, X. Lai, D.W. Goodman, J. Phys. Chem. B 104
(2000) 3050.
[134] K. Luo, D.Y. Kim, D.W. Goodman, J. Mol. Catal. A 167 (2001) 191.
[135] J.A. Rodriguez, M. Kuhn, J. Horbek, J. Phys. Chem. 100 674 (1996)
18240.
[136] A. Tanaka, Y. Takeda, M. Imamura, S. Sato, Phys. Rev. B 68 (2003)
195415.
[137] J. Radnik, C. Mohr, P. Claus, Phys. Chem. Chem. Phys. 5 680 (2003)
172.
[138] P. Zhang, T.K. Sham, Phys. Rev. Lett. 90 (2003) 245502.
[139] G.K. Wetheim, S.B. DiCenzo, S.E. Youngquist, Phys. Rev. Lett, 687 51
(1983) 2310.
[140] B. Richter, H. Kuhlenbeck, J.-J. Freund, P. Bagus, Phys. Rev. Lett. 691,
93 (2004) 026805.
[141] Kim, Y. D.; Wei, T.; Wendt, S.; Goodman, D. W. Langmuir (2003), 19,
7929-7932.
[142] Suzuki, S.; Bower, C.; Watanabe, Y.; Zhou, O. Appl. Phys. Lett. (2000),
76, 4007-4009.
[143] R. Campbell, J. Rodriguez, D.W. Goodman, Surf. Sci. 240 (1990) 71.
[144] V. Vijayakrishman, A. Chainani, D.D. Sarma, C.N.R. Rao, J. Phys.
Chem. 96 (1992) 8679.
[145] K.-H. Meiwes-Broer et al; Metal Cluster at Surface, Springer, (2000).
[146] P. Jena, S. N. Khanna, B. K. Rao; Cluster and Nanostructure interfaces,
World Scientific, 1999.
[147] U. Kreibig and W. Vollmer, Optical Properties of Metal Clusters,
Springer Series, Materials Science, Vol. 25 (1995).
C References
141
[148] Alvarez, M. M.; Khoury, J. T.; Schaaff, T. G.; Shafigullin, M.
N.;Vezmar, I.; Whetten, R. L. J. Phys. Chem. B (1997), 101, 3706-3712.
[149] Vinod, C. P.; Kulkarni, G. U.; Rao, C. N. R. Chem. Phys. Lett. (1998),
289, 329-333.
[150] Beasley, M. R., et al, Phys. Rev. Lett, (1989), 63, 672.
[151] W. Goodman, et al., Phys. Rev. Lett (1989), 62, 2180.
[152] Cordes, O.; Harsdirff, M. Appl. Surf. Sci. (1988), 33, 152-159.
[153] Costanzo, E.; Faraci, G.; Pennisi, A.R.; Ravesi, S.; Terrasi, A.;
Margaritondo, G., Solid State Communication (1992) 81, 155-158.
[154] Wagner, C. D. Anal. Chem. (1972), 44, 967-973.
[155] Wagner, C. D. Anal. Chem. (1975), 47, 1201-1203.
[156] Kleiman, G. G.; Landers, R.; Nascente, P. A. P.; de Castro, S., Phys.
Rev. B (1992), 46, 4405.
[157] Kleiman, G. G.; Landers, R.; Nascente, P. A. P.; de Castro, S,. Surf. Sci.
(1993), 287/288, 798.
[158] Taylor, K. J.; Pettiette-Hall, C. L.; Cheshnovsky, O.; Smalley, R.E. J.
Chem. Phys. (1992), 96, 3319-3329.
[159] S. A. Nepijko et al., Langmuir 15, 5309 (1999)
[160] Luo, K., Lai, X, yi, C. –W. Davis, K. A. Gath, K. K. goodman D. W. J.
Phys. Chem.. B (2005), 109, 4064-4068
[161] H. Hövel, Th. Becker, A. Bettac, B. Reihl, M. Tschudy, E.J. Williams, J.
Appl. Phys. 81 (1997) 154.
[162] S. Zafeiratos, S. Kennou, Surf. Sci. 443 (1999) 238.
[163] W. Eberhardt, P. Fayet, D.M. Cox, Z. Fu, A. Kaldor, R. Sherwood, D.
Sondericker, Phys. Rev. Lett. 64 (1990) 780.
[164] K.H. Hövel, B. Grimm, M. Pollman, B. Reihl, Eur. Phys. J. D. 9 (1999)
595.
[165] M.S. Chen, D.W. Goodman, Catalysis Today, 111 (2006) 22.
[166] D. W. Goodman, Encyclopedia of Nanoscience and Nanotechnology,
Eds. J. A.
142
C References
[167] H.-G.. Boyen, A. Ethirajan, G. Kästle, F. Weigl, P. Ziemann, G..
Schmid, M.G. Garnier, M. Büttner, P. Oelhafen, Phys. Rev. Lett. 94
(2005) 016804.
[168] M. Chen, Y. Cai, Z. Yan, D.W. Goodman, J. Am. Chem. Soc. in press
(2006)
[169] M. Kuhn, T.K. Sham, Phys. Rev. B. 49 (1994) 1647.
[170] A. Bzowski, T.K. Sham, R.E. Watson, M. Weinert, Phys. Rev. B. 51
(1995) 9979.
[171] José A. Rodriguez and D. W. Goodman, J. Phys. Chem. (1991), 95,
4196- 4206
[172] A. Christensen, P. Stoltze, and J. K. Norskov, J. Phys.Condens. Matter
7, 1047 (1995).
[173] M. P. Andrews and S. C. O’Brien, J. Phys. Chem. 96, 8233 (1992).
[174] P. P. Edwards, R. L. Johnston, and C. N. R. Rao; Metal Clusters in
Chemistry, Wiley-VCH, Weinheim, Vol. 3, p. 1454. (1999).
[175] Sinfelt, J. H., Bimetallic Catalysist; Wiley: New York, (1983).
[176] Campbell, C. T. Annu. Rev. Phys. Chem. 1990, 41, 775.
[177] D. C. Lim, Dissertation in proceeding about Chemical Behaviour of Ag
and Au Nanoparticles and Clusters, Universität Konstanz (2006-2007).
[178] D. C. Lim, R. Dietsche, M. Bubek, G. Ganteför, Y. D. Kim:
ChemPhysChem (2006); 7(9) 1909-11.
[179] L. Guczi, G. Lu, Z. Zsoldos, Catal. Today 17 (1993) L103.
[180] M. Baumer, H-J. Freund, Angew. Chem. Int. Ed. 41 (2002) 4073.
[181] M. P. Mallin, C. J. Murphy, Nano Lett 2 (2002) 1235.
[182] T. Shibata, B. A. Bunker, Z. Zhang, D. Meisel, C. F. Vademan, J. D.
Gezelter, J. Am. Chem. Soc. 124 (2004).
Acknowledgment
First, I would like to thank my thesis advisor Dr. Young Dok Kim for his
efficient and intelligent support and for his leadership in our
“Nachwuchsgruppe” that made this thesis possible. In addition, this work would
have never been possible if it were not for the long-time (more than 4 years)
supervision provided by my “Doktorvater” Prof. Gerd Ganteför. Beyond
scientific stuff, I have learned from him to appreciate how important a pleasant
environment inside a working group can be. I would like to express my
gratitude to Prof. Leiderer for his disposition to be the second referee of my
dissertation.
I thank Dong Chan Lim and Rainer Dietsche for the experiments carried
out in the XPS labour, our successful team-work and for their disposition to
help me with every problem during the writing of the thesis.
Other really great thanks go to Moritz Bubek, Tim Fischer, Matthias Götz
Tobias Mangler who have read the first draft versions of some of the chapters
of this thesis. Specially, I wish to thank Jörn Cordes and Marco Niemietz for
the time spent by correcting my written English. It must be hard to read some of
the sentences that I wrote in the first versions of this work.
I am very grateful to my colleges and ex-colleges of AG Ganteför for the
cooperative spirit and the excellent working atmosphere. Interesting discussions
during the coffee breaks, group meetings and the tradition of bringing some
cakes to celebrate “anything” are nice traditions. You are great colleagues and I
am very grateful to all of you!
I also want to acknowledge the staff of the Fachbereich Physik from the
University Konstanz for his disposition to help me at any time.
Finally, I wish to express my grateful to my friends and family, who
support me at any time.