CO oxidation on metal supported
ultrathin oxide films
vorgelegt von
Diplom-Chemieingenieurin
Yulia Martynova
aus Krasnokamensk
Von der Fakultät II - Mathematik und Naturwissenschaften
der Technischen Universität Berlin
zur Erlangung des akademischen Grades
Doktor der Naturwissenschaften
Dr. rer. nat.
genehmigte Dissertation
Promotionsausschuss:
Vorsitzender: Prof. Dr. rer. nat. Michael Gradzielski
Gutachter: Prof. Dr. rer. nat. Hans-Joachim Freund
Gutachter: Prof. Dr. rer. nat. Reinhard Schomäcker
Tag der wissenschaftlichen Aussprache: 30. Mai 2013
Berlin 2013
D 83
Diese Arbeit wurde im Zeitraum von Februar 2010 bis Februar 2013 am Fritz-Haber-Institut der
Max-Planck-Gesellschaft, Berlin, Deutschland, Abteilung Chemische Physik, unter der Leitung
von Herrn Prof. Hans-Joachim Freund angefertigt.
Acknowledgements
I am very glad to express my sincere and deep acknowledgement to my supervisor,
Prof. Dr. Hans-Joachim Freund, who kindly granted me the great opportunity to
perform the high-level research on the leading edge of science and shared his priceless wisdom.
I was happy to gain new knowledge and work with excellent international professionals in the
Department of Chemical Physics of the Fritz Haber Institute of the Max Planck Society.
I am very thankful to Prof. Dr. Reinhard Schomäcker (the Technical University of Berlin)
for kindly being the co-supervisor and reviewing my thesis. As well I gratefully acknowledge
that Prof. Dr. Michael Gradzielski (the Technical University of Berlin) agreed to be the chairman
of the defense examination committee.
I am heartily grateful to Dr. Shamil Shaikhutdinov, my scientific advisor on a daily basis,
for his constant guidance, help and support. I highly appreciate his mastership and experience.
As well I am obliged a lot to Dr. Ying-Na Sun for her patient teaching and introducing me in the
world of surface science.
I am really indebted to my colleagues Dr. Bing Yang, Xin Yu, Dr. Jorge Anibal
Boscoboinik, Dr. Daniel Löffler, Dr. Irene M.N. Groot, Bo-Hong Liu, Martin E. McBriarty and
Emre Emmez for their contribution to my work and versatile help. Furthermore, it was my
pleasure to collaborate with Jonas Weissenrieder and Markus Soldemo (Royal Institute of
Technology KTH, Sweden). I cordially thank Uwe Härtel, Burkhard Kell, Matthias Naschitzki,
Walter Wachsmann, Klaus-Peter Vogelgesang, Max Schönberg and Daniel Jüterbock for their
qualified technical support and assistance. I am grateful to Manuela Misch and Gabriele
Mehnert for their irreplaceable help in administrative deals. I would like to thank as well all
other members of the Department, past and present, for providing help with all kinds of things I
would not have managed on my own.
Especially I want to express my gratitude to my mother and friends for their
encouragement and confidence in me. They inspired me to be hard-working and grafted
optimism.
Abstract
Thin oxide films supported on metal substrates were studied in low temperature CO
oxidation at near-atmospheric pressures. A representative selection of ultrathin films includes
ruthenium oxide films grown on Ru(0001) and Pt(111) and manganese and zinc oxide films
prepared on Pt(111). The performed experimental research has established relationships
between the films structure and their reactivity under technologically relevant conditions (a
particular dependence in each case). Based on our own findings, a unifying concept defining
catalytic behavior of thin oxide films as catalytic materials was proposed. The results suggest to
use oxygen binding energy on the film surface as a good descriptor for the reactivity. Weakly
bound oxygen species can be formed on the oxide film surface at the high chemical potential of
oxygen and serve as active species, which then replenished from the gas phase. Long-range film
ordering and film/substrate nature influence the reactivity to the extent they provide weakly
bound oxygen.
The fundamental study comprises surface science approach for the films
preparation/characterization and conventional reactivity measurements. The films were grown
either by deposition of a foreign transition metal oxide on a supporting metal single crystal or
by oxidation of the substrate material under well-controlled conditions in ultrahigh vacuum.
The film structure and composition were verified by low energy electron diffraction and Auger
electron spectroscopy respectively. The film coverage as well as different oxygen states were
distinguished by temperature programmed desorption. The catalytic performance of the films
(and supports) was quantified via gas chromatography.
This systematic research reduces the lack of interpretation and comparison of the thin
oxide films reactivity. The novelty of the present work relies on the idea that it lifts up thin film
catalysis on the systematic level. The findings of this thesis can aid in rational design of
oxidation catalysts.
Zusammenfassung
Ultradünne Oxidfilme auf Metallsubstraten wurden in der CO-Oxidation bei niedrigen
Temperaturen und quasi-atmosphärischen Drücken untersucht. Eine repräsentative Auswahl
ultradünner Filme enthält die Rutheniumoxidfilme gewachsen auf Ru(0001) und Pt(111) und
Mangan- und Zinkoxidfilme aufbereitet auf Pt(111). Die durchgeführte experimentelle
Forschung hat die Struktur-Reaktivitätsbeziehungen der Filme unter technologisch relevanten
Bedienungen eingerichtet (individuell in jedem Fall). Ausgehend von unseren Ergebnissen
wurde ein vereinheitlichtes Konzept vorgeschlagen, das die katalytischen Eigenschaften von
dünnen Oxidfilmen beschreibt. Die Ergebnisse legen nahe, dass die Bindungsenergie von
Sauerstoff auf der Filmoberfläche als Maß für die Reaktivität verwendet werden kann. Schwach
gebundene Sauerstoffspezies können auf der Oxidfilmoberfläche unter dem hohen chemischen
Sauerstoffpotenzial gebildet werden und als aktive Spezies dienen, die dann aus der Gasphase
ersetzt werden. Die Fernordnung der Filme und die Art der Filme/Substrate beeinflussen die
Reaktivität in dem Maß, in dem sie schwach gebundenen Sauerstoff bieten.
Die Grundlagenforschung verbindet so genannte 'surface science approach' zu
Filmpräparation und Charakterisierung mit herkömmlichen kinetischen Messungen. Die Filme
wurden entweder durch der Ablagerung der fremden Übergangsmetalloxide auf dem
Einzelkristall des stützendes Metals oder durch Oxidation des Substratmaterials unter den gut
kontrollierten Bedienungen im Ultrahoch-Vakuum hergestellt. Die Filmstruktur und
Zusammensetzung wurden anhand der 'low energy electron diffraction' (LEED) und der
Augerelektronenspektroskopie (AES) überprüft. Die Filmbedeckung sowie die verschiedenen
Sauerstoffzustände wurden mittels temperatur-programmierter Desorption (TPD)
unterschieden. Das katalytische Verhalten der Filme (auch der Substrate) wurde mithilfe
Gaschromatographie quantifiziert.
Diese systematische Untersuchung vermindert die Lücken der Interpretation und des
Vergleiches der Reaktivitäten dünner Oxidfilme. Die Neuigkeit dieser Arbeit vertraut auf die
Idee, dass sie die Dünnfilme-Katalyse auf das systematische Niveau hebt. Die Ergebnisse dieser
Arbeit können den rationalen Entwurf von Oxidationskatalysatoren fördern.
Contents
1 Introduction.....................................................................................................................1
2 Background and Methods...............................................................................................5
2.1 Apparatus....................................................................................................................5
2.2 Analytical methods......................................................................................................9
2.2.1 Low Energy Electron Diffraction (LEED)..............................................................9
2.2.2 Auger Electron Spectroscopy (AES)...................................................................15
2.2.3 Temperature Programmed Desorption (TPD)...................................................19
2.2.4 Gas Chromatography (GC).................................................................................30
3 CO oxidation over ruthenium oxide films on Ru(0001).................................................38
3.1 Catalytic activity of Ru-based materials. ...................................................................38
3.2 Structure of materials and thin film preparation.......................................................40
3.3 Results and discussion................................................................................................43
3.3.1 Thickness dependence.......................................................................................43
3.3.2 Active phase formation......................................................................................46
3.3.3 Activation energy and reaction orders...............................................................49
3.3.4 Surface order influence......................................................................................52
3.3.5 Comparative study of RuOx films on Pt(111) vs RuOx/Ru(0001)........................54
3.4 Summary.....................................................................................................................59
4 CO oxidation over zinc oxide films on Pt(111)................................................................60
4.1 Catalytic activity of ZnO-based materials and growth of ZnO films. ..........................60
4.2 Structure of materials and thin film preparation.........................................................63
4.3 Results and discussion..................................................................................................69
4.3.1 Thickness dependence.........................................................................................69
4.3.2 Active phase formation and structural stability of ZnO films..............................74
4.4 Summary......................................................................................................................80
5 CO oxidation over manganese oxide films on Pt(111).....................................................81
5.1 Catalytic activity of MnOx-based materials and growth of MnOx films. ......................81
5.2 Structure of materials and thin film preparation.........................................................84
5.3 Results and discussion..................................................................................................89
5.3.1 Active phase formation and special requirements for its existence...................89
5.3.2. Oxygen states differentiation..............................................................................93
5.4 Summary.......................................................................................................................97
6 General trends in CO oxidation over thin oxide films and concluding remarks…….……..98
6.1 General trends in CO oxidation over thin oxide films……………………………………………….98
6.2 Concluding remarks………………………………………………………………………………………………..101
Bibliography.........................................................................................................................103
List of abbreviations…………………………………………………………………………………………………...…..117
List of figures and tables…………………………………………………………………….…………….……..…..….118
Appendix A: List of publications and conferences attended………………………………….…...……124
Appendix B: Curriculum Vitae........………………………………………………………………………..…...…..126
Erklärung………………………………………………………………………………………………………….…….....……127
1
Chapter 1
Introduction
Production of valuable chemicals always stays an important area of mankind
interests: industry producing materials cares about economic prosperity, while scientists
look for elegant fundamental solutions and analytical methods. To govern a chemical
reaction, i.e. to get a desired product with reasonable time and financial expenses and
suppress byproducts formation or to decompose harmful substance, catalysts are used. This
wide class of materials makes possible catalysis, a phenomenon of chemical reactions
acceleration. Among a broad range of catalytic reactions a special place belongs to those
which occur on the interface between different phases, i.e. heterogeneously. Commonly
heterogeneous catalysis refers to gaseous components interaction on the solid surface. This
is determined by advantageous properties of solid catalysts, since they satisfy to such critical
requirements as high activity and selectivity, prolonged lifetime and easy regeneration [1].
Essential part of conventional catalysts often represents metals dispersed over oxide
support. Fundamental studies needed for precise investigation of the reaction mechanism (a
sequence of elementary steps) demand well-controlled conditions for catalysts preparation
and powerful analytical tools for their characterization. Such a combination is fulfilled in
ultrahigh vacuum (UHV), when catalysts are least affected by foreign molecules, and
electron based surface sensitive techniques can be applied [2]. Under UHV conditions model
(two-dimensional, 2D-) catalysts are simplified to single crystalline surfaces with an ordered
structure and well-known composition. The epoch of intensive studying of surface reactions
of small molecules, e.g. CO oxidation (well-known probe reaction) [3], was facilitated by
advances in vacuum technique and yielded in a special interdisciplinary area of human
knowledge known as surface science [4].
A model reaction of CO oxidation was chosen as a 'working horse' of surface science
to study catalytic properties of metal and oxide surfaces on fundamental level and avoid a
wide range of intermediate reaction products. Such an approach minimizes the overall
complexity of experiments and calculations both. This leads to better understanding of
elementary steps of the reaction and is usually called as 'model studies in heterogeneous
2
catalysis'. On metals the reaction proceeds through the Langmuir-Hinshelwood mechanism
when both reactants are adsorbed on the catalyst surface [5]. On oxygen-covered metallic
surfaces and oxides the Eley-Rideal mechanism could be in operation: only one of the
reactants (O) is chemisorbed, while the other species (CO) reacts by direct collision from the
gas phase [6].
However, studies of oxide single crystals often suffer from charging problems due to
their insulating properties. To overcome this drawback single crystalline thin oxide films,
epitaxially grown on metal single crystals, were introduced into the practice of model
heterogeneous catalysis and serve as supports for deposited metal nanoparticles [7]. But,
this decrease in material dimensions and reactant pressures caused specific problems called
material and pressure gaps respectively: particular material properties and catalytic behavior
depend on catalyst size and pressure scale and cannot be directly extrapolated [8]. Under
the term 'pressure gap' it is commonly understood that the structure of the active phase of a
catalyst on atomic scale under industrial reaction conditions can be completely different
from that in UHV due to several orders of magnitude (≥ 12) pressure drop of the reactants
[9]. To bridge the pressure gap between the fundamental surface science studies and large-
scale chemical processes it was suggested to combine preparation of single crystalline
surfaces and their analysis in UHV with high-pressure reactivity measurements - up to 1 bar
(see [10] and references therein]).The material gap implies that the number of active sites in
industrial catalysts is higher (2-4 orders of magnitude) than in planar (2D) single crystalline
model catalysts, and their quality (coordination, defects, phases) can differ substantially [8].
Recent studies have shown that ultrathin oxide films on metals can demonstrate
interesting catalytic properties on their own right [11-13]. Particularly, in the earlier works of
our group well-ordered monolayer FeO films on Pt(111) were demonstrated to catalyze the
CO oxidation reaction much more active than either bare Pt(111) substrate or nm-thick
Fe3O4 films at low temperatures and near-atmospheric pressures [14]. Such enhanced
reactivity was attributed to the charge transfer between metal and oxide which leads to
oxygen activation on the surface and consequent oxide lattice reorganization. O-terminated
O-Fe-O trilayer structure forming under reaction conditions was proven to direct CO
oxidation through the Mars-van Krevelen like mechanism: weakly bound top layer oxygen
3
species are attacked by gas phase CO resulting in CO2 production and are later replenished
from the gas phase O2 (rate limiting step) [15-17]. For the structural explanation see Fig.1.1a.
Figure 1.1 Ultrathin FeO film grown on Pt(111): a - Lattice transformation of FeO film to FeO2 structure
which supplies weakly bound oxygen for the reaction [11], b - supported Pt nanoparticles are encapsulated by a
FeO layer upon annealing [15].
Therefore the film thickness may affect the reaction, particularly when charge
transfer is the rate limiting step. If the film thickness does matter, then this allows one to
tune the reactivity and selectivity of such an “inverted” catalyst, where an oxide phase is
supported by a metal, in contrast to a traditional oxide-supported metal catalyst. In fact,
such ideas have already been put forward, most notably by Vol’kenstein [18] and Schwab
[19] (see also a critical review by Slinkin et al. [20]), who developed the so-called “electron
theory of catalysis” based, in essence, on the concept of Schottky barriers. However, in that
time the preparation and atomic scale characterization of thin oxide films was not feasible.
In terms of this theory one can explain low reactivity of thick Fe3O4 films: the oxide film
should be very thin to allow the charge transfer, since thin films possess lattice flexibility
which may significantly contribute to stabilization of charged adsorbates (polaronic
distortion). This property is not present in a thicker film, which behaves like the bulk
counterpart. However, high reactivity of Pt nanoparticles deposited on top of thick Fe3O4
films was rationalized as strong metal support interaction (SMSI effect): upon annealing Pt
surface gets covered with a FeO monolayer as shown on Fig.1.1b [21, 22]. Minimization of
surface energy is considered as one of the main driving forces for encapsulation.
It is getting obvious, that this experience with accumulated knowledge is already a
solid foundation in the field of supported oxide film catalysis. At the moment, this fast
developing experimental trend (and theoretical as well) suffers from the lack of descriptive
functions and 'methodological algorithms' which are needed to compare/predict the
a
b
4
reactivity of a new kind of model catalysts. Therefore, it is necessary to give a
straightforward clue and build up an appropriate concept to be able correctly to interpret
the catalytic activity of thin oxide films.
Being inspired with the enhanced reactivity of ultrathin oxide films, mentioned
above, in the present work I have studied the catalytic performance of ruthenium oxide films
grown on ruthenium and platinum and zinc and manganese oxide films grown on platinum.
The aim of this thesis was to establish a direct relationship between the films structure and
their reactivity in CO oxidation with a particular focus on thickness influence.
For this I have prepared a series of thin oxide films of each type under well-controlled
conditions. The films were characterized (order, composition, coverage/thickness) under
UHV conditions and tested in low temperature CO oxidation at elevated pressures. The work
is organized as follows. At first, the whole experimental setup and utilized methods are
overviewed (Chapter 2): low energy electron diffraction (LEED), Auger electron spectroscopy
(AES), temperature programmed desorption (TPD) and gas chromatography (GC). Such a
combined approach of traditional surface science tools (LEED, AES and TPD) and widely used
analytical technique of GC let us to perform accurate film fabrication, their precise
characterization and the reactivity measurements on the high level. Major experimental
results (LEED patterns, AES and TPD spectra, the reaction kinetics) for each kind of films are
presented: for RuOx/Ru(0001) and RuOx/Pt(111) in Chapter 3, for ZnO/Pt(111) in Chapter 4
and for MnOx/Pt(111) in Chapter 5. Since the films nature is rather different, a literature
overview for each film type is given in the beginning of a corresponding chapter prior to
experimental results. In Chapter 6 I compare the catalytic performance of all films and finally
summarize overall findings to conclude about the reactivity of ultrathin films in general and
give an outlook in this respect.
5
Chapter 2
Background and Methods
2.1 Apparatus
A prior requirement for working with the well-defined surfaces is to maintain them in
clean state within analysis time limits. Therefore to minimize the influence of undesirable
adsorbate species (contamination, reconstruction) the samples should be kept under
ultrahigh vacuum (UHV) conditions with pressures lower than 10-9 mbar. Another major
reason concerns a prerequisite for surface science instrumentation. The use of
characterization tools needs a UHV regime as well to avoid electron scattering by residual
gas and to extend the lifetime of an electron source. Therefore a proper vacuum technique is
a necessary precondition to prepare and study model catalysts.
The experiments were carried out in a stainless steel UHV chamber, as shown in
Fig.2.1, with the base pressure 1×10-10 mbar. The vacuum of several ranges is maintained by
four pumping stages which comprise an Edwards rotary pump (~ 10-2 mbar), a Pfeiffer
turbomolecular pump (10-8 mbar) and a Varian Ti sublimation pump with a Varian ion pump
(10-10 mbar). The chamber is heated up to 145⁰C (‘baked’) for at least 12 hours to remove
water vapor as a major contaminant and thus to reach the UHV level. The pressure in the
chamber is monitored by two gauges: in the range 10-6 ÷ 10-10 mbar by a Varian nude
Bayard-Alpert type ionization gauge, in the range 10-5 ÷ 10-4 mbar by a Varian cold cathode
ionization gauge. Gas line system allows supplying six different gases of high purity for
backfilling, high pressure or desorption experiments and is pumped out by an Edwards
rotary pump (~ 10-2 mbar) as well.
The rotatable x-y-z manipulator with an option of liquid nitrogen cooling is employed
to place the sample attached to the sample holder in front of various analytical tools. The
metal substrate used in this work is a double-side polished single crystal (99.99%, MaTeck,
~10 mm in diameter, 1.5 mm in thickness) of either Pt(111) or Ru(0001). As shown in Fig. 2.2
the sample was mounted to two parallel Ta (1mm) sticks by a pair of Ta wires (0.25-0.4 mm),
6
which were used for resistive heating and also for cooling by filling a manipulator rod with
liquid nitrogen. The temperature was measured by type K thermocouples (chromel-alumel)
spot-welded to the edge of the crystal, and controlled by a PID feedback system (Schlichting
Phys. Instrum.). The temperature could be varied between ca. 85 K and 1300 K.
Figure 2.1 Horizontal cross-sectional view of the UHV chamber
For the thin films deposition by the physical vapor deposition (PVD) method the UHV
chamber was equipped with electron beam assisted metal evaporator (Focus EFM3T) and a
homebuilt effusion cell containing high purity deposition materials (Goodfellow). Clean
metallic surfaces were prepared by Ar+ ion bombardment with an Omicron sputter gun.
Low energy electron diffraction (LEED) optics coupled with an Auger electron
spectrometer (AES) from Omicron was used to control the surface quality of prepared
samples. A Hidden quadrupole mass spectrometer (QMS) was utilized for Residual Gas
Analysis and Temperature Programmed Desorption (TPD) studies. For TPD studies the
sample was exposed to gases through a directional gas doser. The pressure in that line was
controlled by a MKS Baratron® gauge (MKS Instruments) and maintained by a Pfeiffer
7
turbomolecular pump. The LEED patterns were recorded with a Canon EOS 400D camera.
The AES and TPD data collected by the corresponding software were evaluated using Origin
program (OriginLab).
Figure 2.2 Sample holder and high-pressure cell sketch
The reactivity studies and treatments at elevated pressures were conducted in an
especially designed “high-pressure” cell (HP cell, ~30 ml, Au-plated Cu massive block), built
in inside the UHV chamber (see Fig. 2.2). An Agilent Technologies gas chromatograph (GC)
connected by gas lines to the HP cell allows monitoring gas composition and the reaction
kinetics. The manipulator rod inside the chamber ends with a KF-type flange with a 4-pins
electrical feedthrough holding Cu and thermocouple wires.
To carry out the CO oxidation reaction, the following procedures were used. The
well-characterized sample was moved vertically down to the high-pressure cell, which was
then sealed from the UHV chamber by using a Viton O-ring placed on top of the reactor
8
matching the flange at the end of the manipulator. The reactor was then filled with reactants
(CO (99.997%) and O2 (99.9996%), Linde) up to the desired pressure in mbar range and
balanced by He (99.9999%, Linde) up to 1 bar. Gas admission was controlled by two MKS
Baratron® gauge (MKS Instruments). Meanwhile the main chamber remained in the pressure
of ~10-8 mbar range. CO was additionally cleaned using a cold trap kept at 200 K. Reactants
were circulated for 20 min at room temperature to equilibrate gas concentrations and to
steady the gas flow using a membrane pump. The sample was then heated up to the
reaction temperature with a constant heating rate. Sampled gas mixture probes were
divided in GC using an HP-PLOT Q column in isothermal mode at 35 ºC and analyzed by a
thermal conductivity detector (TCD) (see Fig. 2.3).
After the reaction, the spent catalyst was cooled down to room temperature within
2-3 min, and the reactor was pumped out down to ~10-5 mbar before exposing to UHV. Then
the manipulator was retracted, and the post-reaction structural characterization of model
catalysts was performed.
Figure 2.3 Gas chromatograph setup
9
2.2 Analytical methods
2.2.1 Low energy electron diffraction (LEED)
LEED stays the most frequently used experimental technique for structural
characterization in surface science. Upon bombardment of crystalline solids by accumulated
beam of electrons with a primary energy below 300 eV and observation of diffracted
electrons as spots on a fluorescent screen the method can provide the information about
long-range ordering and symmetry [23]. Due to different interatomic and interplanar
distances in crystalline materials, the electron diffraction patterns are characteristic for a
specific crystallographic structure and can be used to distinguish the structures.
Figure 2.4 Diagram of a standard 4-grid LEED optics apparatus
A typical experimental set up is illustrated schematically in Fig. 2.4. The standard
LEED apparatus comprises a hemispherical fluorescent screen and an electron gun aligned
along the central axis of the screen. Accelerated electrons are emitted from the electron gun
filament with low work functions (LaB6, W or Ir), which then follow paths along the central
axis with the aid of Wehnelt cap and electrostatic lenses, and finally hit the sample at ground
potential in order to prevent charging. The backscattered electrons are then accelerated
10
towards the screen biased to 3-6 kV. The grid nearest the sample is grounded to allow the
linear beam trajectories. The next two grids (called retard, suppressor or gate) are at a
potential of several volts below the incident electron beam and provide accurate energy
selection, therefore the inelastically scattered electrons which have lost more than a few
volts cannot reach the screen – they are screened out. To make the retarding field
homogeneous and mechanically stable these grids often go in pair. The fourth grid is again
grounded to prevent the influence from the screen potential on the energy-selection grids.
The elastically backscattered from crystal atoms electrons give a diffraction pattern [24].
In 1921 L. de Broglie predicted that every moving particle has a wave-like nature. This
means that the particles, e.g. the electrons, undergo the effects associated with waves, such
as diffraction or interference. In 1927 C. Davisson and L. H. Germer experimentally observed
the angular distribution (a diffraction pattern) of electrons elastically scattered from Ni
sample when applying a monochromatic beam of low energy electrons onto the crystal
surface and explained their data in terms of diffraction of the electrons from crystallites [24].
The theoretical explanation for the formation of diffraction patterns from crystal
structures is as follows: coming from de Broglie’s equation for the electron’s wavelength, the
wavelength’s dependence on the electron’s kinetic energy (determined by the electron
beam energy) can be derived:
√
√ (
)
⁄ (2.1)
where λ is the wavelength in m, h=6.626×10−34 Js is the Planck’s constant, p is the electron’s
momentum, Ek is the electron’s kinetic energy, m=9.11×10−31 kg is the electron mass,
e=1.6022×10−19 C is the electron charge and V is the beam energy.
Electron scattering provides information about the surface (is surface sensitive) if the
electrons are granted the correctly chosen energy. The inelastic mean free path (IMFP) of
electrons in a solid depends on the energy of the electrons in a manner that does not
depend too strongly on the chemical identity of the solid illustrated by the universal curve
(see Fig.2.5) [25]. A beam of low energy electrons (V = 30-200 eV) results in a small
wavelength λ of 1-2 Å which is of a size of interatomic distances in a crystal lattice (< 3 Å).
11
This is short enough to probe only first 3-4 atomic layers: electrons scattered from deeper
gap virtually play no role in the detected signal.
Figure 2.5 Universal curve of electron mean free path in solid matter [25]
Diffraction spots experimentally visible on a display system certify that the elastically
scattered electrons undergo a constructive interference following the Bragg’s law (for X-rays
originally) if the path difference between the electrons coming from different atoms is equal
to the multiplicity of the wavelength [24]:
(2.2)
where d is the interplanar distance in the crystal lattice, Θ is the angle between the beam
and the sample plane, n is an integer and λ is the wavelength. The process of electron
scattering at the crystal surface is schematically shown in Fig. 2.6. The electron beam used at
above 50 eV has typically a beam diameter of around 1 mm. The instruments transfer width
(which is essentially the coherence length of the electrons) is of approximately 5-10 nm in
modern instruments. Naturally only structures which extent over at least (100 x 100) Å can
produce Bragg reflexes. Structure domains with a smaller diameter than the coherence
width will not constructively add to the diffraction pattern but add diffuse background
intensity. The observed diffraction pattern is therefore an ensemble average of structural
domains (larger than the incident beam coherence width) within the electron beam.
12
Figure 2.6 Diffraction process
LEED gives the information about the 2D-structure of the sample surface. A
diffraction pattern is a geometrical visualization of a real crystal lattice in reciprocal space –
the concept well-known for bulk X-ray diffraction. In 2D space, the lattice vectors are
spanned by two base vectors and , the reciprocal lattice follows from the definition of
the reciprocal lattice vectors and
:
,
(2.3)
is the vector normal to the surface. This also means:
(2.4)
| |
|
| , |
|
|
| (2.5)
where are the base vectors of the real lattice ( =1,2); are the base vectors of the
reciprocal lattice (𝑗=1,2); is the Kronecker delta, = =1, , = =0; α is the enclosed
angle between and . The diffraction conditions for a two dimensional lattice are given
by the two dimensional Laue conditions:
(
) , (
) (2.6)
where h and k are arbitrary integers. They principally represent the conservation of energy
and momentum between the incident wave vector
and the emerging wave vector
. This
condition is fulfilled whenever
equals a reciprocal lattice vector
. The
13
energy conservation requires that |
| |
|. These two conditions can be made visible by
changing the Ewald construction known from x-ray scattering to the surface case as shown in
Fig. 2.7 [26].
Figure 2.7 Ewald sphere synthesis for 2D case. Incident and diffraction beams are labeled
An overlayer structure with a rather large unit mesh and various domains may
generate a complicated pattern. The lattice vectors of the surface including possible
adsorbate overlayers are characterized as
and
. A simple nomenclature of surface
structures is that of Wood’s. The surface structure is described by
(
) (2.7)
where N= “P” or “c” for primitive or centered cells, respectively, and is the angle by which
the surface vectors have to be rotated with respect to those of the bulk (see Fig. 2.8). The
nomenclature of Wood has the advantage of simplicity. It is, however, not possible to
describe all surface structures because the rotation angle might not be the same for both
vectors.
14
Figure 2.8 Terminology for surface lattices. 𝒂𝟏 and 𝒂𝟐 correspond to the substrate unit cell vectors. 𝒃𝟏 and 𝒃𝟐
indicate the superstructure unit cell vectors. 𝜶 is the enclosed angle between the substrate, 𝜸 is the enclosed
angle between the superstructure unit cell vectors. The angles 𝜷 and 𝜹 accounts for a rotation of the
superstructure with respect to the substrate.
The matrix representation (offered by Park and Madden) of superstructures shows that as
and
can be built up by the substrate unit cell vectors:
,
(2.8)
[
] (2.9)
The indices 𝑗 can be derived by trigonometric considerations:
[ | |
| |[ ]| |
| |
| |
| |[ ]| |
| |
] (2.10)
where the meaning of the variables is illustrated in Fig. 2.8.
In addition to mathematical formalism of the method theory it is necessary to label
the real structures of the crystal surfaces and superstructures resulting from coincidence
with adsorbate overlayers as well as corresponding LEED patterns using fundamental
crystallographic two-dimensional (2D) Bravais lattices and Miller indices for low index planes
as shown on Fig.2.9.
15
Figure 2.9 Crystallographic notation systems. A: 2D Bravais lattices: a – square, b – rectangular, c –
centered rectangular (rhombic), d – hexagonal, e - oblique. ). B: Miller indices. C: Real space- and reciprocal
lattices (white spots are associated with the adsorbate structure.
2.2.2 Auger electron spectroscopy (AES)
AES is a widely used experimental technique utilizing the Auger effect. The technique
allows determining surface chemical composition by investigating the kinetic energy of the
Auger electrons emitted from the sample upon electron bombardment. One of practical
advantages of the AES technique is experimental compatibility with LEED optics. The Auger
effect was discovered by Lise Meitner and Pierre Auger independently in the 1920s [27]. It
was first described by Meitner; however, Pierre Auger was a scientist who was mostly
credited for it. They observed an internal conversion of energy from a radiationless
transition.
(2x1)
(100)
A
C
B
16
Figure 2.10 Auger electron emission mechanism and shells nomenclature
The physical mechanism of the Auger process and the shells nomenclature are given
in Fig. 2.10. An electron beam of high energy typically of 2-5 KeV is used to ionize core levels
of atoms, labeled as A. The ionized atom that remains after the removal of the core electron
is in a highly excited state (a) and will rapidly relax back to a lower energy state by one of
two routes: X-ray fluorescence (process b) and Auger emission (process c). A higher level
electron fills the vacant core level. The energy liberated in this process is transferred to a
second electron which is called the Auger electron. The Auger electron is then emitted with
the kinetic energy:
Ek = Ecore - Eouter1 - Eouter2 - δE - φ (2.11)
where the δE term accounts for the relaxation effects involved in the decay process, which
leads to a final state consisting of a heavily excited, doubly ionized atom; φ is the work
function.
Therefore, the kinetic energy of the Auger electron is independent of the primary energy. An
Auger transition (ABC) is therefore described by the initial hole location (A) and followed by
the locations of the final two holes (B, C). The energies of the Auger electrons are in the
range of 50 to 3000 eV and give a characteristic spectrum of peaks for atoms of a specific
element due to unique orbital energies (besides H and He no Auger peaks because of two-
17
hole final state). Since AES peaks are rather low in intensity (see Fig.2.11), the spectra are
typically recorded as the first derivative using lock-in techniques as shown in Fig. 2.12 [28].
Figure 2.11 Energy spectrum of electrons coming from a surface irradiated with a beam of primary electrons.
Electrons have lost energy to vibrations and electronic transitions (loss electrons), to collective excitations of
the electron sea (plasmon losses), and to all kinds of inelastic processes (secondary electrons). The element-
specific Auger electrons appear as small peaks on an intense background, and are better visible in a derivative
spectrum.
Figure 2.12 AES spectrum of ruthenium oxide as an example. Primary energy: 3 keV.
The AES often bases on the same setup as LEED utilizing a retarding field analyzer
(see Fig.2.4 and Fig.2.13). Electron gun is placed in front of the sample and current is applied
to the filament. Emitted and accelerated electrons hit the sample initiating the Auger
process. Retarding grids allow only the Auger electrons with a given energy to pass through.
The potential increases in steps, for example 1 eV/1 sec, which allows collecting all electrons
with energies in a given range. The multiplier filters noise, strengthens and digitizes the
signal which is then recorded by a computer in a form of the kinetic energy vs. Auger
electron intensity plot. It is possible by adding a small sine (periodic oscillation) to the gate: a
phase-sensitive detection yields an output signal proportional to the derivative.
18
Figure 2.13 AES - electron energy analysis
Although Ek is independent of the primary beam energy (Ep), Ep is important for
obtaining optimum Auger yield. The ionization probability Qi to form a core hole in a level of
energy Ei depends on the energy of the core level as 1/ Ei2, as well as on the ratio Ep /Ei . In
order to remove an electron with a binding energy Ei one needs a primary electron of higher
energy. If the primary energy is too high, the electron simply goes too fast to have efficient
interaction with a bound electron in a core level. As a result, it is shown theoretically by
calculating cross section that a maximum cross section is obtained when Ep is three to five
times of the core electron energy (see Fig.2.14) [28]. Therefore, the Auger electron of the
kinetic energy ranging from 50 eV to 1000 eV need the primary energy in the range 2-5 keV.
Auger transitions are more probable for lighter elements. According to the universal curve
shown in Fig. 2.5, the Auger electron can escape from subsurface region of 1-2 nm in depth
which corresponds to 3-5 atomic layers of the sample.
Figure 2.14 Optimal energy of a primary beam and yield of Auger electrons. a - The probability Qi to create a
core hole in a level with binding energy Ei with a primary electron of energy Ep maximizes for Ep/Ei ≈ 2–3. b -
Auger decay (solid curve) is the preferred mode of de-excitation in light elements, while X-ray fluorescence
(dotted curve) becomes more important for heavier elements (shown for K shell vacancies, similar plots can be
obtained for L and M shell transitions).
19
A quantitative analysis of Auger spectra is easily possible, but less straightforward
than in X-ray photoemission spectroscopy (XPS) due to backscattering effects [28]. For
example, the Auger yield measured from a few atomic layers thick molybdenum film on a
tungsten substrate is about 20% higher than that of bulk molybdenum, because the heavier
tungsten is an efficient backscatterer. This effect should be taken into account by multiplying
the primary electron intensity ip by a factor of (1+R), where R is the backscattering factor. In
general, R is higher for heavier than for lighter elements, but it may also be matrix
dependent. The general expression for the Auger electron current ia becomes:
∫
Where Y is the probability of Auger decay; n(z) is the concentration of the element at depth
z; λ is the inelastic mean free path of the Auger electron; Θ is the take-off angle of the Auger
electron measure from the surface normal. For a homogeneous sample with the detector
perpendicular to the sample surface, Eq.2.12 reduces to
The mole fraction of component A in a binary mixture of A and B is given by
⁄
⁄⁄
where IA and IB are the intensities of the Auger peaks and sA and sB are the relative sensitivity
factors. If there are more than two components then the denominator has to be replaced by
a sum over all components. The mole fraction can be equated to the relative coverage [29].
2.2.3 Temperature-Programmed Desorption (TPD)
TPD is a surface science technique used for determination of kinetic and
thermodynamic parameters of desorption processes or decomposition reactions from
desorption behavior of adsorbed species. It can provide both qualitative and quantitative
information. A sample exposed to a gas is heated with a temperature program (where the
temperature is a linear function of the time) and the partial pressures of atoms and
molecules evolving from the sample are measured, e.g. by mass spectrometry. When
20
experiments are performed using well-defined surfaces of single-crystalline samples in a
continuously pumped UHV chamber then this experimental technique is also referred to as
thermal desorption spectroscopy (TDS). TPD offers interesting opportunities to interpret
desorption in terms of reaction kinetic theories, such as the transition state formalism.
The sketch of a typical experimental setup for TPD studies is presented on Fig.2.15.
The sample is linearly heated with a rate β=0.5-5 Ks-1 resulting with the temperature ramp
T=T0+βt with time t. The concentration of desorbing species is usually measured with a
quadrupole mass spectrometer (QMS) also known as a residual gas analyzer (RGA). The
pumping speed should be sufficiently high to prevent readsorption of the desorbed species
back onto the surface. The mass spectrometer has a shield with a nozzle pointing towards
the sample which therefore eliminates signals coming from the sample holder. The sample is
positioned < 1 mm far from the nozzle. A QMS consists of three sections. In the first
ionization region, gas atoms and molecules are ionized by electrons (typically of ~70 eV)
produced by W cathode. Ions are then accelerated and focused into the second section, the
mass filter, which consists of four parallel bars forming an electrical quadrupole field. The
applied voltages affect the trajectory of ions traveling down the flight path centered
between the four rods. For given voltages, only ions of a certain mass-to-charge ratio pass
through the quadrupole filter and all other ions are thrown out of their original path. A mass
spectrum is obtained by monitoring the ions passing through the quadrupole filter as the
voltages on the rods are varied. The last section of the spectrometer is the ion detector
which contains a secondary electron multiplier [28].
Figure 2.15 Schematic representation of QMS setup for TPD experiment
21
Among the range of surface processes which can occur on gas-solid interface such
phenomena as adsorption and desorption are important for understanding of interatomic
matter interaction. In surface science adsorption is implied as a process of concentration of
species from a gas phase on the surface of a solid. The rate of adsorption is governed by the
rate of impingement of molecules at the surface and the sticking coefficient. This flux per
unit surface area (F, molecules·m-2·s-1) is given by the Hertz-Knudsen equation:
√ (2.15)
where P is gas pressure, N·m-2, k is Boltzmann's constant (1.38×10-23 K·J-1), m is molecular
mass of the impinging species, kg and T is the absolute temperature. The rate of adsorption
(Rads) is
(2.16)
Depending on the nature of the interaction-potential there is a certain probability that the
atom/molecule is not immediately scattered back into the gas-phase but comes to reside at
the surface. The probability of this event is referred to as the sticking coefficient (S) and
takes a value from 0 to 1 where the value of 1 means that all atoms/molecules impinging the
surface spend significantly longer time at the surface compared to in the direct scattering
process. It is essential here to mention that S depends on the substrate temperature as well
as on the temperature of the incoming molecule, but also it strongly depends on the number
of available adsorption sites at the surface, decreasing with decreasing number of available
adsorption sites. In the case of high availability of surface adsorption sites, and low
temperatures at which adsorbate can form condensed multilayers, one can assume S to be
essentially constant.
Adsorption of atoms or molecules from the gas phase onto a solid surface occurs
when attractive forces exist at short distance between them. There are two principle modes
of adsorption on surfaces: physisorption and chemisorption. Physisorption is a weak
bonding, with typically low enthalpy values 20-40 kJ·mol-1, characterized by the lack of a true
chemical bond between adsorbate and surface, i.e. no electrons are shared, not saying that
there cannot be significant electron rearrangement in the molecule upon adsorption. So, the
22
basis of distinction is the nature of the bonding between atoms or molecules and surfaces.
The weakest interaction which can lead to bonding between a surface and an adsorbate is
the attractive van der Waals interaction, which originates from the interaction of a
fluctuating dipole in the adsorbate with a polarizable surface. A van der Waals bonding
between two molecules can be described as the interaction between two point dipoles. In
the case of bonding to a surface, the attractive potential has a r−3 dependence. However, as
the atom is brought very close to the surface, the overlapping of the electron cloud of the
adsorbate and the substrate leads to a steep increase of kinetic energy of the electrons and
hence a high repulsive potential as a result of the Pauli exclusion principle. Thus a
physisorption potential is the sum of the repulsive and attractive Van der Waals
contributions [30] (see Fig. 2.16, potential energy diagrams).
Chemisorption includes also all other types of interactions resulting in a stronger net
chemical bonding to the surface compared to physisorption. Chemisorption occurs when
there is an overlap between the electronic orbitals of the adsorbate and the surface, leading
to the formation of chemical bonds of energies typically exceeding 50 kJ·mol-1. The bond can
be predominantly ionic or covalent. There are three types of chemisorption: molecular
chemisorption (a), non-activated dissociative chemisorption (b) and activated dissociative
chemisorption (c) (see Fig. 2.16) [27]. The potential energy curve is described by a
combination of the physisorption and chemisorption. The left region from the crossing point
in the figure is dominated by the chemisorption. The depth of the adsorption well is a
measure of the strength of the binding to the surface and is represented as adsorption
energy (𝐸 ). The desorption of an adsorbed species from a surface is considered to be the
reverse process of adsorption. For cases (a) and (b), the desorption energy (𝐸 𝑒 ) is equal to
𝐸 . The dissociative chemisorption is often an activated process, which means that an
activation barrier (𝐸 𝑐𝑡) needs to be overcome. 𝐸 𝑒 then can be equated:
𝐸 |𝐸 | |𝐸 | (2.16)
The potential energy diagram for an activated dissociative chemisorption process is shown
on Fig. 2.16(c). A diatomic molecule approaching the surface from a large distance first
physisorbs on the surface. If a molecule possesses kinetic energy enough to overcome the
activation barrier, the molecule may dissociate. The dissociated species can be bonded to
23
the surface in a chemisorption state. The corresponding potential curve of two atoms differs
from that of the molecule at the large distance by the dissociation energy of 𝐸 .
Figure 2.16 Potential energy diagrams for a diatomic molecule approaching and interacting with a
solid surface
The energetics and kinetics of adsorption can be understood by evaluation of the TPD
data [31-33]. Fig. 2.17 schematically shows a typical TPD spectrum (pressure-temperature
curve). The profile of desorption is characterized by a desorption peak - maximal
concentration of desorbing species at certain temperature in agreement with the Boltzmann
distribution. Each peak corresponds to various desorbing species. The area under a peak is
proportional to the amount of adsorbed species. The position of the peak (the peak
temperature) is related to the enthalpy of adsorption. TPD provides information on the
strength of the bond between adsorbate and substrate.
Figure 2.17 Schematic TPD spectrum
24
The experimental aspects of thermal desorption spectroscopy are based on the
concepts of adsorption and desorption. If a solid with a surface free of adsorbed particles is
placed in a gas medium, the process of adsorption will commence. Under static conditions,
this process will continue in static until an equilibrium concentration of adsorbed particles is
reached. If the solid surface under consideration is placed inside in a chamber of volume (V ),
which is under a constant evacuation at speed S, and the gas is leaked at a constant rate (L),
the above equilibrium concentration can be described in terms of the steady-state pressure
(peq) inside the chamber:
(2.17)
Here K represents the proportionality constant (K = 3.27×1019 molecules·l-1, at p = 1 torr and
T = 295 K). If the sample surface is heated at a constant rate, the gas pressure inside the
chamber increases due to the flux of desorbing particles. Assuming no re-adsorption of the
desorbing particles, the mass balance of particles leaving and entering the chamber is given
by adsorption and desorption processes:
(2.18)
where A is the surface area of the adsorbent, dN/dt is the rate of desorption and p is the
instantaneous pressure inside the chamber. Adsorption on the walls of the system has also
been neglected. Using Eq. 2.17, and substituting Δp=p-peq, the equation of mass balance [Eq.
2.18] can be simplified as
(2.19)
Here a=A/KV and =V/S. is the characteristic pumping time. Depending on the magnitude
of the pumping speed, two experimental scenarios can be identified. At a small pumping
speed ( →∞), the desorption rate is proportional to the first derivative of pressure with time
(dN/dt ~ dΔp/dt). Alternatively (for the most part of TPD studies), at high pumping speed
( →0), desorption rate is proportional to the pressure (dN/dt ~ Δp).
25
Analysis of desorption signals is based on treating the desorption process as a kinetic
phenomenon. The kinetics of the desorption process follows the rate law with Arrhenius
type behavior described as:
𝑒 (
) (2.20)
where R=8.314 J·K-1·mol-1 is the universal gas constant, T is the absolute temperature in K, Θ
is the coverage (0≤Θ≤1), ν is the pre-exponential frequency factor (in many cases = 1013 s-1
is a good approximation), n is the desorption order, and 𝐸 is the desorption activation
energy, kJ·mol-1. A frequently used basis for the evaluation of kinetic parameters from TPD
spectra is known as the Polanyi-Wigner equation defining the rate of desorption Rdes:
𝑒 (
) (2.21)
here ß is heating rate in K·s-1. Initially the increasing of the temperature results in an
exponential rise of the desorption rate. At the same time coverage of the surface decreases.
The parameters 𝐸 , n, and ν should generally be considered as functions of coverage and
may in cases of complex kinetics (e.g., a phase equilibrium on the surface) also depend on
desorption temperature, which is a function of the heating rate. That means that desorption
parameters can depend on the conditions of the TPD experiment.
The surface coverage Θ is a function of the gas pressure over the sample. It is
characterized be the Langmuir adsorption isotherm which is widely used for the discussion
of adsorption/desorption phenomena [34]. The concept is based on the following
assumptions: a - adsorption is localized, i.e., the adsorbed particles are immobile; b - the
substrate surface is saturated at Θ = 1 ML (monolayer), i.e., when all adsorption sites are
occupied, c - there are no interactions between the adsorbed particles.
The desorption order, if interpreted as the molecularity of an elementary reaction,
allows insight into the nature of the rate-limiting step of the desorption process, which is,
from the kinetic point of view, in general a consecutive reaction. From the molecularity of
the desorption process one can often draw conclusions about the state of the adsorbate
itself. Zero-order desorption (n=0) takes place (see Fig.2.18 a), if the desorption rate does
26
not depend on the coverage. The spectra are typically asymmetric and show all the same
leading edge, while the temperature of maximum desorption shifts to higher values with
increasing initial coverage. First-order desorption (n=1) spectra have symmetric line shapes
(Fig.2.18 b), the peak temperature of which is independent of the initial coverage. Second-
order desorption (n = 2) is typical for dissociative adsorption (see Fig.2.18 c). In this case, the
peaks are symmetric, and the temperature at maximum desorption shifts to lower values
with increasing initial coverage.
Figure 2.18 Orders of desorption process. Sets of TPD profiles illustrate various desorption kinetics: (a)
- zero order, (b) - first order, (c) - second order.
In the literature, several methods for evaluating the desorption activation energies,
prefactors and orders of desorption have been described, including mostly Redhead’s
analysis, leading edge analysis and complete analysis.
In 1962 Redhead [33] derived a simple relation between the desorption activation
energy 𝐸 , the peak maximum temperature Tmax and the desorption order assuming that
activation parameters are independent of surface coverage and that desorption followed
first order kinetics. The maximum in the desorption rate occurs when Rdes/ 𝑇=0. Redhead
has shown that the desorption rate equation for =1 can be solved to obtain:
𝑒 (
) (2.22)
Solving Eq.2.22 for 𝐸 provides:
𝐸 𝑇 (
) (2.23)
a
b
c
27
The second term in brackets is small relative to the first and is estimated as 3.64. The error
introduced through this estimate is less than 1.5% for 108 < / < 1013 K-1. The Redhead
method is often employed to extract activation energies from a single peak desorption
spectrum.
Leading edge analysis was introduced by Habenschaden and Küppers in 1984 [31].
According to this method it is expected that in this small temperature range (~ 4% of
spectra), the variation of ν can be neglected, and is close to the initial coverage. The
leading edge of the TPD curve fits the Polanyi-Wigner equation (Eq. 2.21). The desorption
rate is proportional to the intensity of the TPD trace. Here the equation becomes:
𝑒 (
) (2.24)
where is proportionality constant. Taking logarithms of both sides of this equation leads
to:
(2.25)
If is then plotted against 1/T, from the slop of such Arrhenius plot the value of 𝐸
can be determined. From the intercept with the ordinate the frequency factor can be
determined provided that reaction order and coverage are known. If is plotted
against at a fixed temperature, the slop will be n, order of desorption. The great
advantage of this method is that it requires a minimum of assumptions. A drawback is the
use of the leading edge for the analysis because the signal/noise ratio in this region of low
desorption rates is inherently low.
The first complete analysis of TPD was given in 1975 by King [32]. Fig. 2.19 illustrates
the procedure: a coverage is chosen (e.g. 0.15 ML); the points corresponding to =0.15
on all TPD curves are found. This gives a pair of (Rdes and 𝑇) values from every curve with
initial coverage larger than 0.15 ML. An Arrhenius plot of all versus 1/T values for this
particular coverage yields 𝐸 . The prefactor follows from the intercept + , when
the order of desorption is known. For coverage above 0.1 ML the term is much smaller
than , and could be ignored. The pre-exponential factor can be regarded as representing
28
the frequency of attempts of the adsorbed particle to escape the chemisorptive potential.
The values determined vary by at least four orders of magnitude, from 1012 to 1016 s-1.
Figure 2.19 Complete analysis of TPD data. (a) The spectra are integrated to determine points on the
spectra corresponding to a fixed coverage. (b) A pair of ( , T) values for every desorption trace are given,
from which an Arrhenius plot is made. (c) The slope yields the activation energy [28]
The transition state theory of reaction rates gives the link between macroscopic
reaction rates and molecular properties of the reactants, such as translational, vibrational
and rotational degrees of freedom [28]. The desorption of a molecule M proceeds (see
Fig.2.20):
→
→
(2.26)
where is the adsorbed molecule;
is the molecule in the transition state for
desorption; is the equilibrium constant for the excitation of into the transition
state; is the Boltzmann’s constant; h is the Planck’s constant; T is the temperature; is
the molecule desorbed in the gas phase.
29
Figure 2.20 Desorption process in terms of the transition state theory
For desorption, the reaction coordinate is the vibration of the molecule with respect to the
substrate. The chance that the adsorption bond breaks is given by the factor 𝑇/ . Therefore
the rate constant of desorption in the transition state theory is:
𝑒 (
)
(2.27)
where is the partition function of
, is the partition function of , 𝐸 is the
adsorption energy. The partition function contains translational, vibrational and rotational
terms:
(2.28)
√
;
;
(2.29)
where is the characteristic linear dimension of an adsorption site; m is the mass of the
molecule M. is the vibrational frequency; 𝐼 is the moment of inertia of the molecule M. The
total partition functions are the products of each term for each individual translation,
vibration and rotation. The simpliest case arises if the partition functions and are
almost equal. This corresponds to a transition state that resembles the ground state of the
adsorbed molecule. The activation energy 𝐸 𝑐𝑡 equals to 𝐸ads + 𝑇 for desorption through a
tight transition state. According to the definition of the activation energy:
𝐸 𝑇
(2.30)
30
Comparison with Eq. 2.27 results in the prefactors equal to
, 1.6×1013 −1 at 300 K.
Higher pre-exponential factors are obtained when the molecule rotates or moves in the
transition state, indicating rather loose transition state, for example the prefactor for
desorption of CO from Ru(0001) is in the range of 1015 to 1016 s-1.
2.2.4 Gas chromatography (GC)
GC represents a widely applied analytical tool for chemical compounds identification
in a mixture based on their separation between the mobile gaseous phase and the stationary
phase. The method is practically realized in a gas chromatograph (GC).
The operational principle and major parts of the GC involved in the experimental
procedure are schematically presented on Fig.2.21 (also see general overview on Fig. 2.3).
Upon manual or automated sampling by the injector the analyte mixture is introduced in the
GC, vaporized in a heater if necessary and diluted by a carrier gas (mobile phase). Then the
gas flow by means of a rotary valve enters the column, a tube filled (packed) or coated
(capillary) with an immobile material, where the chromatographic separation takes place
through a distribution of components due to molecular interaction with the column
material. The column is thermostated in the oven, a section used for the temperature
variation. Further divided substances in the order of separation follow to the detector where
they are finally analyzed utilizing specific physical properties. An electronic signal produced
by the detector is plotted against time by the computer assisted software as a
chromatogram. The latter consists of a number of peaks corresponding to components in
the sample. The peak integrated area shows the concentration of a certain substance while
the peak position (time) refers to its chemical nature [35].
Distribution of a component is determined by its relative affinity for the two phases.
Those solutes interacting more strongly with the stationary phase will be retained longer in
the column. Molecular interaction results from intermolecular forces that are electrical in
nature. In GC such types termed as dispersion forces and polar forces are important. All
interactions between molecules are composites of these forces. Dispersion forces arise from
charge fluctuations throughout a molecule resulting from electron/nuclei vibrations. Polar
31
interactions arise from electrical forces between localized charges resulting from permanent
or induced dipoles.
Figure 2.21 Operational principle of a general purpose gas chromatograph. Pneumatic control is not shown.
The mathematical background for quantitative analysis in GC relies on the basics of
plate and rate theories. The concept of the plate theory was originally proposed for the
performance of distillation columns. However, Martin and Synge (1941) first applied the
plate theory to partition chromatography. The theory assumes that the column is divided
into a number of zones called theoretical plates. One determines the zone thickness or
height equivalent to a theoretical plate (HETP) by assuming that there is perfect equilibrium
between the gas and liquid phases within each plate. The resulting behavior of the plate
column is calculated on the assumption that the distribution coefficient remains unaffected
by the presence of other solutes and that the distribution isotherm is linear. The diffusion of
solute in the mobile phase from one plate to another is also neglected [36].
Eluting from the column the component is characterized by the (absolute) retention
time 𝑡 - the amount of time that elapsed from injection of the sample to the recording of
the peak maximum of the component band (peak) by the detector. This term is commonly
used as the adjusted retention time 𝑡
- the solute total elution time minus the retention
time for an unretained peak of the mobile phase (or holdup time -the time necessary for the
32
carrier gas to travel from the point of injection to the detector; this is characteristic of the
instrument, the mobile phase flowrate, and the column in use):
𝑡
𝑡 𝑡 (2.31)
The column efficiency or sharpness of the peak is defined by the theoretical plate number N.
Assuming the peak to have normal (Gaussian) distribution (see Fig. 2.22) N could be found:
(
) (
) (
) (2.32)
where is the peak band broadening, is the peak width at base and is
the peak width at the half of the height.
Figure 2.22 Schematic drawing of a component peak on a chromatogram.
Having calculated the number of theoretical plates and knowing the length L of the column,
one may determine the HETP (or simply H):
𝐸𝑇
(
) (2.33)
Thus Eq.2.33 shows that the HETP represents the peak broadening as a function of retention
time. In a gas chromatographic column, each component will yield different N and HETP
33
values. Those solutes with high retention will result in greater numbers of theoretical plates
and thus lower HETP values.
Peaks resolution defined as a degree of separation between neighbour peaks of two
components 1 and 2 in terms of their average peak widths can be written as follow:
(2.34)
Capacity (retention) factors k is a measure of the ability of the column to retain a sample
component. When an analyte retention factor is less than one, elution is so fast that
accurate determination of the retention time is very difficult. High retention factors (greater
than 20) mean that elution takes a very long time. Ideally, the retention factor for an analyte
is between one and five. Selective components distribution is characterized by the selectivity
factor . It is in turn expressed through the ratio of k for two components 1 and 2:
(2.35)
The separation factor is the ratio of the distribution constants K for substances 1 and 2
measured under identical conditions. By convention the separation factor is usually greater
than unity:
(2.36)
where is a component concentration in the stationary phase and its concentration in
the mobile phase. Finally, taking into account capacity and separation factors the resolution
could be given by a commonly used expression as follows:
(
)(
) (2.37)
where N and k refer to the later-eluting compound of the pair. If two peaks are separated by
a distance , then Rs = 1. If the peaks are separated by a , then Rs = 1.5 which practically
means the beginning of an effective separation. If further to consider the non-negligible time
for substances in the mobile phase, this gives the effective theoretical plates related to the
number of theoretical plates:
34
(
) (2.38)
Although HETP is a useful concept, it is empirical. One must use the rate theory to
explain chromatographic behavior. The rate theory is based on such parameters as rate of
mass transfer between stationary and mobile phases, diffusion rate of solute along the
column, carrier-gas flowrate, and the hydrodynamics of the mobile phase. The van Deemter
equation (1956) is used for describing the gas chromatographic process. The equation was
derived from consideration of the resistance to mass transfer between the two phases as
arising from diffusion:
(2.39)
where A - eddy diffusion term, B - longitudinal or ordinary diffusion term, C - or resistance to
mass transfer term and u is a linear gas velocity. A representation of this equation is given in
Fig. 2.23, which shows the effect of H with changes in linear gas velocity. This equation
predicts that for maximum column performance, the contribution of each term must be
minimized while maintaining a constant linear flowrate.
In the case of a capillary column, the A term is equal to zero because there is no
packing material. Thus, Eq. 2.39 simplifies to abbreviated expression often referred to as the
Golay equation (1957):
(2.40)
Figure 2.23 Van Deemter plot. Change in H versus linear gas velocity u
35
In the practice of chromatographic instrumentation the central role is played by
columns. They are categorized into two classes by their operational principle (see Fig. 2.24).
The first type originally introduced is packed one when the separation of the gas mixture
takes place in a column filled with particles of a solid occupying the major volume. In
chemical analysis packed columns were replaced by capillary columns in the early 1980s with
mass production of fused silica tubes. Inner walls of a capillary column with the narrow
diameter are coated with a thin layer (μm) of non-volatile immobilized liquid or solid
substance as a stationary phase. This offered a tremendous improvement in resolving power
compared to a conventional packed column. Such a column is also often referred to as an
open tubular column. The high permeability or low resistance to carrier-gas flow of capillary
column enables a very lengthy column to generate a large number of theoretical plates.
Figure 2.24 Representative flow through chromatographic columns: a- a packed column, b - a capillary column
The proper selection of a stationary phase/column depends on the nature of mixed
substances to be separated. Generally, one can be guided by the rule of matching polarity:
polar components require a polar stationary phase material and vice versa. Among the
general purpose polar materials polysiloxanes, polyethylene glycol are widely used for
various organics analysis. Graphitized carbon and porous polymers belong to non-polar
adsorbents and are applied in hydrocarbons separation. Molecular sieves divide well noble
and permanent gases.
The optimal chromatographic performance of the column is defined by combination
of its inner diameter, length, the carrier gas flow (linear velocity), the stationary phase
36
thickness and the temperature mode. The larger diameter/flow/higher temperature
facilitates shorter analysis times but results in lower peaks resolution and backwards. The
length change influences in a reverse way: a longer column will have higher resolution
compensated by the increased time of detection. Temperature-programmed operation is
preferred rather than isothermal when separating components with close physico-chemical
properties (molar mass, boiling point). A thicker layer of the stationary phase film causes the
magnification of the column capacity (mass of retained analyte) and the retention time as
well. Finding the balance of all parameters depends on the particular experimental problem
(and the column price) [36].
To register the chemical substances a wide variety of detectors is nowadays applied
in GC technique, often coupled with other methods. The detector performance is
characterized by a number of important parameters. The limit of detection (LOD) is the
concentration at which the output is equivalent to twice the noise level. The dynamic range
of a detector is usually given as a concentration ratio to LOD within which a concentration
dependent output is produced. The most significant region of the dynamic range is the linear
range. An ideal gas chromatographic detector should have low LOD and a large linear
dynamic range (five to six orders of magnitude). The response factor is the ratio of the signal
(integrated area)-to-sample size and must be stable within the linear dynamic range. A
selective detector responds to compounds containing a specific heteroatom while a
universal detector responds to any component. Detectors can be divided into two groups:
mass flow detectors, which respond to the mass of sample component reaching the detector
per unit time, and concentration-sensitive detectors, which provide an output that is directly
proportional to the concentration of a sample component in the mobile phase [37].
The most popular detector in GC practice stays a thermal conductivity detector (TCD
or katharometer) due to its cost and accuracy. The TCD detector responds to any substance
different from the carrier gas. The detector (see Fig.2.25) contains two filaments: one
exposed only to carrier gas, while the other is exposed to the carrier gas for sample analysis.
Instead of a direct measurement of filament temperature, the filament resistant, which is a
function of temperature, is measured. When the gas for the sample analysis is only carrier
gas, the two filaments can be balanced. Such a scheme is called the Wheatstone bridge. The
37
ability of a colliding molecule to carry off heat depends on its thermal conductivity.
Hydrogen and helium have high thermal conductivity and therefore will be more efficient at
“cooling” a heated filament than other gases. When the sample cell filament is subjected not
only to the carrier gas the primary signal is produced [36]:
(2.41)
where is a resistance difference between two filaments, α is the temperature coefficient
of resistance for the filament wire, the resistance at the reference temperature of 0⁰C, λ
is proportional to the temperature difference across the cell, T1 is the temperature of the
filament wire, T2 is the temperature of the detector.
Figure 2.25 The working principle of a thermal conductivity detector. R3 and R4 are the resistance of the sample
and reference cells respectively, R1 and R2 are the auxiliary resistances.
Qualitative data analysis i.e. identification of substance by the peak position (the
retention time) is carried out by comparison with that of pure standards or test mixtures for
the same experimental conditions and column parameters. Quantitatively the amount of
each component in a mixture is found as:
∑
(2.42)
where is integrated area of the component peak, is the response factor of the
detector for this substance (it could be found from a calibration curve).
38
Chapter 3
CO oxidation over ruthenium oxide films on Ru(0001)
3.1 Catalytic activity of Ru-based materials.
Ru(0001) is famous to interact with a broad range of versatile small molecules (often
coadsorbed) like CO, O2, H2, NO, N2, NH3, H2O, hydrocarbons, methanol [38-49] and to be
active in such chemical processes as CO oxidation [50], ammonia synthesis and oxidation
[51, 52] under UHV conditions. Its surface can also serve as a substrate for preparation of a
wide variety of thin films. This includes oxides of transition metals (MOX, where M=Ce, Ti, Fe,
Al), silica, graphene and bimetallic alloys as ordered layers [53-59].
However, its native oxide coating the ruthenium bulk continues to attract attention
in fundamental research in its own right, especially in the CO oxidation reaction at elevated
pressures. The stoichiometric oxide RuO2 growing in (110) direction on Ru(0001) surface was
found to interact with CO, O2, H2, NO, H2O, hydrocarbons [60-66] and catalyze CO, NH3 and
HCl oxidation [67-69] in UHV. A review of ruthenium dioxide properties and applications in
comparison with metallic ruthenium was recently given by Over [70]. It turned out that CO
oxidation became a very challenging topic in surface science of Ru-based catalytic materials
due to discrepancies and ambiguity raised by the pressure gap and different reaction
conditions what resulted in a number of phases ascribed to be reactive. The milestones of
this intriguing story are provided below.
The beginning of the intensive Ru-related studies takes place at the moment when
the reaction conditions were changed from UHV to mbar pressure range of reactants in the
work of Peden and Goodman where they reported Ru(0001) to be more active than Rh and
Pd in CO oxidation with stoichiometric gas ratio [71]. There they proposed the idea to
explain Ru anomalous behavior and developed it further that a dense (1x1) phase of
chemisorbed oxygen on Ru(0001) is the active phase in the reaction [6, 72].
39
Later Over and co-workers suggested that under technologically relevant conditions
the Ru catalyst is represented by the RuO2(110) surface [73, 74]. Enhanced reactivity was
proposed to be governed by interaction of weakly bound oxygen atoms with CO [75]. Upon
the following-up studies with polycrystalline RuO2 powder and supported Ru nanoparticles
Muhler with colleagues formulated the shell-core model [76, 77]. The group of Schlögl in the
studies with Ru(0001) and polycrystalline RuO2 powder suggested that the active phase is a
thin film of poorly defined oxide structure labeled as “surface oxide” [78, 79].
These investigations induced search of similar oxygen rich surface structures on other
noble metal catalysts at mbar pressures. The formation of ultrathin oxide layers on Pt, Pd,
and Rh and their high reactivity in CO oxidation have recently been reported [80-85].
Experiments were supported by DFT calculations with the full step-by-step analysis of the
oxide formation and CO oxidation energetics over both Ru(0001) and RuO2(110) surfaces by
Reuter and Scheffler [86-89].
Recently, Somorjai with coworkers presented the size-reactivity dependence of CO
oxidation over core-shell Ru nanoparticles covered by a thin layer of oxide species [90].
Therefore, through a controversial interplay between experiment and theory it is now
believed that the active phase is neither a chemisorbed oxygen phase nor the stoichiometric
RuO2(110) but rather a very thin ruthenium oxide film of not yet well determined structure.
Systematic studies of the reactivity of “native” oxide films as a function of the film
thickness have not been performed so far, although in their recent paper Goodman and co-
workers mentioned similarities in the catalytic behavior of the ruthenium oxide films
prepared at different oxidation temperatures [91]. Thus, the efforts to transfer the ideas
introduced in Chapter 1 also to “native” oxide films have a rational background.
Such films could in principle be even formed on noble metal surfaces, provided the
high chemical potential of oxygen. In this work, I have examined thin ruthenium oxide films
on Ru(0001) (as well as bare Ru and chemisorbed O-phases) in the CO oxidation reaction at
near-atmospheric pressures and low temperatures (400 – 470 K). In particular, my study was
focused on finding the relation between the reactivity of oxide films and the film thickness.
40
3.2 Structure of materials and thin film preparation
Ru(0001) has hexagonal close packed (hcp) crystal structure with lattice constant
a=2.704 Å what results in a hexagonal (1x1) LEED pattern. The major surface contaminants
are Si, S and C [92]. Clean surfaces of the Ru(0001) single crystal can be obtained by several
cycles of argon ion sputtering and annealing in UHV and oxygen [93]. However, it is difficult
to confirm the cleanliness of Ru surface since its MNN Auger peak at 273 eV interferes with
the carbon KLL peak at 272 eV. After the surface quality was proved by LEED and AES, it can
be finally verified by CO TPD. CO desorption from clean Ru(0001) is characterized by a well-
known profile with two peaks α1 at ~ 460 K and α2 at ~ 400 K corresponding to undisturbed
√ √ and compressed CO respectively, that was revealed by Menzel with
coworkers [41].
Figure 3.1 Real space geometry of Ru(0001) and O-adlayers unit cells [70]
Several chemisorbed oxygen adlayer structures can be formed on Ru(0001). Their
real space schematic representation is shown of Fig.3.1. The O(2x1)-Ru(0001) surface with
the coverage Θ=0.5 ML was prepared by exposure to 2x10-7 mbar O2 at 420-450 K for 10 min
[39] giving a (2x2) hexagonal LEED pattern which originates from a coincidence of three
rotational domains. Here one ML (monolayer) corresponds to a number of surface atoms
that equals that of Ru atoms in the topmost layer of Ru(0001).
Upon population the surface with O-atoms its sticking coefficient decreases, and
higher oxygen chemical potential is needed to attach more oxygen to the metal. The 3O(2x2)
41
structure with Θ=0.7 ML [94] was prepared by oxidation in 1x10-6 mbar O2 at 1220 K for 5
min and cooling in the same oxygen pressure down to 300 K. The dense O(1x1)-Ru(0001)
surface with O in hcp sites and Θ=1 ML was formed upon the crystal exposure in the HP cell
to 20 mbar of pure O2 at 450 K for 10 min in contrast to 'NO2 method' [95]. The formation of
the oxidic phase under UHV conditions is kinetically hindered [86].
Bulk ruthenium (IV) oxide has a rutile crystal structure where each Ru atom is
completely coordinated by six surrounding oxygen atoms with a unit cell 4.51 Å × 3.11 Å
[96]. But the bulk-truncated surface of new crystallographic plane e.g. (110) is terminated
with rows of bridging oxygens and shows 5-fold Ru atoms called 'coordinatively unsaturated
sites (cus) with missing O neighbours as shown on Fig.3.2. Single crystalline RuO2(110) films
can be grown on the hot Ru metal substrate by dosing high amounts of O2 [73].
Figure 3.2 Surface geometry of RuO2(110) [97, 98]: a - ball-stick model, b - top view on a unit cell, c -
oxide domain on Ru(0001) substrate
In this work native thin oxide films on a double-side polished Ru(0001) crystal were
prepared by oxidation of the crystal surface in 10-4 mbar of O2 at 600 - 700 K by backfilling
a
b
c
42
entire UHV chamber. The film thickness was varied by the oxidation time and temperature.
The freshly prepared oxide films all showed LEED patterns characteristic for the RuO2(110)
overlayer on Ru(0001) as depicted on Fig.3.3d: it evolves from the superposition of a (1x1)
hexagonal arrangement of Ru(0001) support and additional array of RuO2(110)-(1x1) spots
with rectangular symmetry seen as domains rotated 120⁰ with respect to each other; in the
real space the oxide has a unit cell 6.38 Å × 3.11 Å [97, 98].
Figure 3.3 LEED patterns (70 eV) of different O-phases on Ru(0001) used in this work: a - clean Ru(0001), b -
O(2x1) or 3O(2x2), c - O(1x1) and d - RuO2(110). (1×1) and (110) unit cells are indicated.
The surface composition was determined by AES using normalization to O(2×1)- and
3O(2×2)-Ru(0001) structures as the references. The nominal film thickness is presented in
this work in oxygen monolayer equivalents (MLE) such that 1 MLE corresponds to the
amount of the surface oxygen atoms in the (1×1)O-Ru(0001) structure. A typical Auger
spectrum of an oxide film was given as an example on Fig.2.12 in Ch.2. In this work oxide
films were grown up to maximal thickness ~ 7 MLE.
A characteristic CO TPD spectrum from oxide films served as an additional fingerprint
[60, 99]. Based on this data it was assumed that oxide films possess the stoichiometry and
ordering of RuO2(110) phase. However, to prevent any speculations about the film structure
within the whole thickness interval it was decided to use symbolic nomenclature of RuOx.
Additional structural characterization of the similarly prepared oxide surfaces was
performed in another UHV chamber equipped with LEED, scanning tunneling microscopy
(STM) and x-ray photoelectron spectroscopy (XPS) by B. Yang, X. Yu and J.A. Boscoboinik,
members of the 'Structure and reactivity' group, Department of Chemical Physics, Fritz-
Haber-Institute of Max Planck Society.
43
3.3 Results and discussion
3.3.1 Thickness dependence
The CO oxidation reaction was performed in a circulating mixture of CO and O2 in the
mbar-pressure range balanced by He to 1 bar. Typical kinetic curves of CO2 production under
oxidizing conditions (10 mbar CO + 50 mbar O2) are shown in Fig.3.4a. The CO2 production
grows linearly in time during the first 10-15 min, then slows down and stops as CO is totally
consumed. Characterization of the post-reacted surfaces by LEED and AES revealed
practically the same surface composition and ordering, suggesting no substantial structural
transformations under the reaction conditions. For all oxygen adlayer structures, the post-
reacted surfaces showed solely the (1×1) LEED pattern, i.e., in contrast to the high
temperature conditions (>550 K) where the oxide formation is observed during the reaction
[74, 76, 78].
On Fig.3.4b reactivity data are given in terms of turnover frequencies (TOF), the
number of produced CO2 molecules per active sites and second, as calculated from:
𝑇
(3.1)
where V=165.6 cm3 is the volume of the HP cell (reactor) and gas lines to the GC,
L=6.022·1023 mol-1 is Avogadro's number, R=8.413 J·K-1·mol-1 is the universal gas constant,
T=300 K is the gas temperature, is the number of active sites of the catalyst
(the area of the double-side crystal is about 1.08 cm2), P(CO2) is the partial pressure of CO2, t
is time. Since during the reaction the direct measurement of the partial pressure of CO2 was
not provided, it was assumed that its maximal value is equal 10 mbar as well as the initial
pressure of CO, what is based on the full conversion of CO: dP(CO2) ~ dP(CO) ~ 10 mbar.
Figs.3.4 (a, b) clearly shows that thin oxide films are much more active than the clean
Ru(0001) surface. The O-adlayer surfaces showed negligible reactivity, regardless of whether
the surface exhibited O(2x1) or 3O(2x2) structures before the reaction, since the surface
immediately transforms into the (1x1) structure upon exposure to high pressures of oxygen
in the reactor. Each sample was prepared separately and used only once.
44
Figure 3.4 Reaction kinetics for CO oxidation on clean Ru(0001) crystal, O-adlayers and thin oxide films
grown on Ru(0001), possessing the indicated amounts of oxygen (in MLE) prior to the reaction: a - kinetic
curves of CO2 production, b - turnover frequencies of CO2 production.
In order to quantify the effect, the CO2 production rate was measured within the first
10 min, although the oxide films showed substantial conversion, above 50 %. The nearly
constant reaction rate even at high CO conversions implies no poisoning by CO2 that is
a
b
45
accumulated in the circulating mixture. To further prove this conclusion, the following
experiment has been performed as shown in Fig.3.5.
After 10 min of the reaction over a 5 MLE oxide film at 430 K the reactor was pumped
out while cooling the sample to the room temperature. Then, the reactor was re-filled with
the fresh CO + O2 mixture, heated up to the reaction temperature, and the reactivity was
measured again. In the beginning, the CO2 production rate slightly diminished, most likely
because of carbonaceous deposits, but it recovered to the original value as the carbon
burned out in the excess of oxygen.
Indeed, when the same procedure was repeated except the sample has been flashed
to 600 K in UHV prior to refilling the reactor, the reaction proceeds with the same rate from
the onset. Therefore, it was concluded that the reaction does not suffer from CO2 self-
poisoning, at least under net oxidizing conditions. This finding allowed comparing the activity
of oxide films even at high CO conversions.
Figure 3.5 Sequential CO oxidation runs over the same 5 MLE thick RuOx film on Ru(0001) (10 mbar
CO+50 mbar O2, balanced by He to 1 bar, 430 K).
When the reaction rates are re-plotted against the sample oxygen content prior to
the reaction, Fig.3.6 clearly shows a steep increase in the reactivity upon oxygen
incorporation into the Ru surface (i.e. oxygen coverage exceeds 1 MLE). Basically, the CO
oxidation reaction sets in only in the presence of oxide films. Increasing of the nominal film
46
thickness further enhances the reactivity, but to the lower extent. These results suggest that
the presence of a very thin oxide layer is, in principle, sufficient to show superior catalytic
activity, i.e. in contrast to the O-chemisorbed layers which are inactive under the conditions
studied. For clarity a vertical dashed line indicates a point where oxygen content of the film
represents already the stoichiometry of RuO2.
Figure 3.6 Structure-reactivity dependence of CO2 production over RuOx films on Ru(0001),
O/Ru(0001) layers and initially clean Ru(0001) as a function of oxygen coverage measured before the reaction
(10 mbar CO+50 mbar O2, balanced by He to 1 bar, 450 K). The onset of oxide formation is marked by a vertical
dashed line.
3.3.2 Active phase formation
To gain further information about the atomic structure of the films, similarly
prepared films have been examined by STM in another UHV chamber (B. Yang, X. Yu, J.A.
Boscoboinik), where the oxygen coverage was determined by XPS [100].
Relatively thick films (approximately 4 MLE and above) exhibited sharp LEED patterns
of RuO2(110)/Ru(0001) and showed STM images very similar to the previously reported by
Over et al. [73] and Rössler et al. [101], with rectangular-shaped terraces dominating the
large scale morphology (Fig.3.7a). High resolution STM images (inset in Fig.3.7a) revealed the
atomic structure characteristic for the RuO2(110) surface, where the protruding lines with a
47
~ 6.5 Å spacing were assigned to the bridging oxygen atoms in the bulk rutile structure of
RuO2 [97]. At these high oxygen coverages, an oxide film uniformly covered the entire crystal
surface.
Figure 3.7 STM images of RuOx films grown on Ru(0001): (a) Typical large-scale STM image of “thick” (>
4 MLE) films. The inset shows a high-resolution image characteristic for the RuO2(110) surface, with a ~ 6.5 Å
spacing between the protruding rows. (b) STM image of raw-like structures additionally observed on the ultra-
thin films (1-2 MLE). The streaks are caused by the tip instability along the scanning direction. The inset shows
atomically resolved STM image of the rows as well as 3O(2x2)-Ru(0001) structure in between [100]. The images
are provided by B. Yang, X. Yu and J.A. Boscoboinik.
In contrast, the ultra-thin films (i.e., 1-2 MLE), showing faint diffraction spots of
RuO2(110), exhibited very heterogeneous surface, with patches of RuO2(110) coexisting with
the (1x1)O-Ru(0001) surface. In addition, new row-like structures were observed by STM as
shown on Fig.3.7b. The protruding rows were primarily running along the [10-10] direction,
i.e. the same as for RuO2(110) overlayer. However, the rows exhibited a ~4.6 Å periodicity
along the rows (see inset in Fig.3.7b), which is considerably larger than observed for the
bridging oxygen rows on RuO2(110), i.e. 3.1 Å. Also the spacing between the rows (e.g. ~ 9 Å,
on average) is definitely larger than the distance between adjacent rows on RuO2(110) (= 6.4
Å). For the isolated rows it was possible to see by STM that the “open” surface between the
rows exhibits the honeycomb-like 3O(2x2)-Ru(0001) structure (inset in Fig.3.7b).
Tentatively, these rows have been assigned to the one-dimensional oxide structures
as an intermediate state and/or precursor to the formation of RuO2(110) overlayer.
48
Certainly, determination of its atomic structure needs further studies. In this respect the
recent work of Over's group is interesting where it is proposed that the growth of 3D RuO2
clusters requires high O2 pressures only on the initial step of their formation [102]. The
authors claimed no evidence of O-Ru-O trilayer formation predicted by Reuter et al. [87].
Nonetheless, thermal desorption spectra of oxygen on the oxide films of various
thicknesses did not reveal any new features beyond those observed on the “thick”
RuO2(110) films. A set of preformed oxide films was subjected to mbar pressures of pure
oxygen at the reaction temperature. Post-characterization with LEED and AES indicated no
change in surface ordering and composition.
Upon exposure to 10 mbar O2 at 450 K all the films studied showed O2 desorption
signal at ~ 420 K in TPD spectrum as presented on Fig.3.8 (amplified by a factor of 5),
previously assigned to terminal oxygen [61], and the main peak centered at ~ 1010 K
resulted from the film decomposition and thus attributed to lattice oxygen. The intensity of
both peaks basically scaled with increasing the nominal film thickness. Therefore, combined
together the LEED, STM, and TDS results suggest the increase of the reaction rate above 1
MLE, shown in Fig.3.6, due to the increasing the surface fraction covered by an oxidic (not
necessarily the RuO2) phase.
Figure 3.8 TPD spectra of oxygen from decomposition in UHV of oxygen pretreated oxide films
RuOx/Ru(0001) (10 mbar O2 at 450 K, 20 min). Heating rate β=3 Ks-1
49
3.3.3 Activation energy and reaction orders
The apparent activation energy of CO oxidation was measured under oxidizing
conditions only for the 5 MLE samples, where the RuO2(110) film covered the whole surface.
A freshly prepared film of the same composition was prepared for each experiment. Kinetic
curves of CO2 production at different temperatures are shown on Fig.3.9a. In the mixture of
10 mbar CO and 50 mbar O2 the Arrhenius plot in the temperature range of 400 - 470 K
yields the activation energy ca. 58±4 kJ/mol as shown on Fig. 3.9b.
Figure 3.9 CO oxidation at different temperatures over 5 MLE oxide films on Ru(0001) in a mixture of
10 mbar CO and 50 mbar O2 balanced to 1 bar with He: a - kinetic curves, b - the Arrhenius plot.
a
b
50
This value is considerably lower than 78±10 kJ/mol reported for the 1.6 nm
RuO2(110)/Ru(0001) film and measured at the nearly stoichiometric CO/O2 ratios (14 mbar
CO + 5.5 mbar O2) at 470 - 670 K [103]. Since in the current work the reaction was carried
out at low temperatures and excess of oxygen, no deviation from isothermal mode was
observed in contrast to [104], where authors performed the direct comparison between the
reactivity of non-oxidic phase and that of a RuO2(110) layer 1.6 nm thick (250 mbar CO + 130
mbar O2, > 550 K) and reported the ignition of the reaction over the oxidic phase causing
temperature increase up to 130 K and a self-acceleration of the reaction.
The reaction orders n for CO and m for O2 in the reaction rate equation (where r is
reaction rate, k is reaction constant, and surface coverages of CO and O2 respectively)
(3.2)
were determined by measuring reactivity at different CO/O2 ratios at 430 K. In one set of the
experiments, the partial pressure of CO was set to 10 mbar, and the oxygen pressure was
varied. In another set, the oxygen pressure was set to 20 mbar, and the CO pressure was
varied. (Again, for each experiment, a new film with the same surface composition was
prepared, 5 MLE thick). The results, shown in Fig.3.10, revealed the first order reaction for
CO and practically zero order for O2. Zero orders for both CO and O2 were reported by Over
et al. [103], which were, however, determined only qualitatively by varying partial pressures
of both gases simultaneously.
The reaction kinetics, presented in Figs.3.10 (a,b), clearly shows the catalysts
deactivation with time. Several mechanisms for the deactivation have been discussed in the
literature: a surface reconstruction into a less active phase [105] and a carbonaceous (e.g., a
carbonate) contamination [91, 101]. Which of these two is operative in experiments of
current work is difficult to ascertain on the basis of solely AES and LEED characterizations of
the post-reacted surfaces. The results of Fig.3.5 demonstrated no CO2 self-poisoning effect
otherwise expected for the carbonate mechanism. It therefore appears that the deactivation
is a relatively slow process, at least under net oxidizing conditions (see Fig.3.4).
In the recent work of Gao et al. [91] authors demonstrated that under oxidizing
conditions and relatively low temperatures (<450 K) preformed RuO2 films display higher
51
activity than the (1×1)O-Ru(0001) phase, i.e. in full agreement with this work. However, the
activity was determined from the pressure changes monitored with a baratron gauge using
the entire UHV chamber (61.6 l) as the reactor. The reactant mixture was renewed each time
when the conversion exceeded 10 %. But this high reactivity regime for RuO2 is restricted: (1)
to very oxidizing reaction conditions (2) to very low reaction temperatures, and (3) to short
reaction times. Although these are the conditions used in the present work, its results
indicate that the decisive parameter is the low reaction temperature, at which the
(preformed) RuO2(110) films are more active than metallic Ru(0001) regardless of the CO:O2
ratio and reaction time.
Figure 3.10 Pressure dependence of the CO oxidation reaction on 5 MLE RuOx films on Ru(0001) at 430
K: a - the CO pressure was varied at fixed p(O2) = 20 mbar (He balance), b - the O2 pressure was varied at fixed
p(CO)= 10 mbar (He balance), c - reaction rates vs pressures. Each data point corresponds to a new film.
a
b
c
52
It seems the only discrepancy that remains between current work and Gao et al’s
study is their finding of a ‘‘negative’’ activation energy at T<475 K in the mixture of 8 Torr CO
and 40 Torr O2 (see Figs. 7a and 8 in [91]), that has never been reported for any technical or
model Ru catalysts. Meanwhile, regular behavior is shown in the current work, with
activation energy of 58±4 kJmol-1 (see Fig.3.9b). In order to explain the rate decreasing with
the temperature in the range 400–475 K, Gao et al. invoked the formation of a carbonate
which deactivates the active sites at low temperatures, but dissociates upon approaching
500 K. However, this explanation implies that the carbonate overlayer was formed before
the reactivity was measured. In contrast, the experiments shown in Fig.3.5 suggest that the
presence of CO2 (as a precursor for a carbonate) does not change the reaction rate, at least
on the most active films.
3.3.4 Surface order influence
Finally, in order to see the effect of surface ordering on the reaction, the reactivity of
the disordered films was compared with that of ordered films. For this, freshly prepared
oxide films were subjected to mild Ar+-sputtering (500 eV) followed by re-oxidation in 10-4
mbar O2 at 450 K, i.e., much below the temperature used for the preparation of ordered
films (~ 700 K).
These treatments resulted in the disappearance of the characteristic diffraction spots
of the RuO2(110) phase, i.e. loss of its long-range order, and increased background intensity
with Ru spots left on the LEED pattern as shown on Fig.3.11, although no considerable
changes in the surface stoichiometry were observed by AES. Nonetheless, for the two
thicknesses studied (4.5 and 6.5 MLE), the disordered films exhibited higher reaction rate
than the ordered films (see Fig.3.12). Each kinetic curve refers to a new sample. Thermal
decomposition confirmed the presence of O2 peak at ~ 420 K in TPD spectrum similar to that
observed for O-pretreated ordered RuOx films on Fig.3.8.
Thus, the reactivity of oxide surface is not related to the surface ordering, thus
suggesting that CO oxidation over the ruthenium oxide surfaces is, in fact, structure
insensitive. Therefore, the rate enhancement, observed for the disordered films, could, in
principle, be explained by the increased surface area of the roughened surfaces. These
53
results agree well with the previous high pressure (0.1 mbar, CO:O2=1) XPS studies showing
no direct correlation between the high CO2 production rate and the formation of the
stoichiometric RuO2 phase [78].
Figure 3.11 LEED patterns (70 eV) of (a) an ordered 5 MLE RuO2(110)/Ru(0001) and (b) a disordered RuOx film
on Ru(0001) prepared from (a) by 500 eV Ar+ sputtering at 300 K and re-oxidation in 10-4 mbar O2 at 450 K.
Figure 3.12 Effects of surface ordering on the reactivity of thin oxide films on Ru(0001). Solid and
opened symbols show the results for ordered and disordered films, respectively, at two thicknesses as
indicated. The disordered surfaces were prepared by 500 eV Ar+ sputtering at 300 K and re-oxidation in 10-4
mbar O2 at 450 K.
a
b
54
3.3.5 Comparative study of RuOx films on Pt(111) vs RuOx/Ru(0001)
In order to see the effect of a metal substrate thin ruthenium oxide films were grown
on Pt(111). Pt(111) and Ru(0001) planes have the similar geometry, and thus the correct
epitaxial relationships for rutile-like RuO2(110) plane are possible as shown on Fig.3.13
(similar to the sketch on Fig.3.2) as well as for (100) plane [106].
Figure 3.13 Epitaxial relationships of rutile RuO2(110) phase on a hexagonal surface Pt(111) [106]
So far fabrication of thin ruthenium oxide films on Pt(111) under UHV conditions has
never been reported. Mostly the available literature refers to electrochemical growth. In
[107] the authors reviewed the preparation of Ru decorated Pt(111) electrodes by
electrochemical and PVD deposition and its activity in CO electrooxidation attributed to a
surface metallic alloy (with the initial submonolayer Ru coverage). There was briefly
mentioned that the presence of oxide fraction inhibits the ruthenium deposition. Highly
(110)-textured polycrystalline RuO2 films 150-300 nm thick were grown on
Pt(111)/Ti/SiO2/Si(001) by metal-organic chemical vapor deposition (MOCVD) method [108].
300 nm thick RuO2(100)/ Pt(111) electrodes were prepared by reactive sputtering in oxygen
for further use as a substrate [106].
In this work the preparation of thin RuOx films on Pt(111) was made in two steps. At
first ruthenium was evaporated on Pt(111) at room temperature from a metallic rod by e-
beam assisted reactive deposition in 10-7 mbar O2. Subsequent postoxidation in 10-4 mbar O2
was performed at 700 K, identically to the requirements for the growth of the native oxide
film on Ru(0001). As prepared oxide films exhibited the (1×1) hexagonal LEED pattern with
the slightly diffused spots and enhanced background intensity as shown on Fig.3.14b, i.e.
completely different from expected RuO2(110) phase grown on Ru(0001). Upon the film
55
preparation Pt spots should attenuate, however, lattice constants of Pt(111) and Ru(0001)
are similar (2.71 and 2.77 Å respectively) that it is difficult to distinguish in this superimposed
hexagonal array of spots with enhanced intensity between Ru and Pt. Thus, the theoretical
calculation [106] of the epitaxial growth of a rectangular lattice over a hexagonal one did not
come true in the actual experimental work.
Figure 3.14 LEED patterns (70 eV) of RuOx film on Pt(111): a- clean Pt(111), b - RuOx film on Pt(111).
On deposited RuOx films open patches of Pt(111) surface could be excluded, since no
CO uptake typical for Pt(111) qualitatively (by the peak shape) and quantitatively (the peak
position and height) was observed as depicted on Fig.3.15. Desorption of CO from ideal
Pt(111) surface is characterized by well-known broad peak between 300-500 K with a
pronounced shoulder at 360 K [109].
Figure 3.15 CO TPD profiles from clean Pt(111) and RuOx film on Pt(111). Heating rate β=3 Ks-1
a
b
56
The amount of oxygen on the surface clearly indicates the presence of some oxidic
phase (measured both by AES and TPD). The Auger spectra of RuOx/Pt(111) and
RuOx/Ru(0001) look very similar as shown on Fig.3.16.
Figure 3.16 Auger spectra of RuOx films of comparable thickness on Ru(0001) and Pt(111) (clean Pt(111) is given
as well for the reference), 3 keV.
The reactivity measurements indeed demonstrated CO2 formation over
RuOx/Pt(111). The observed production rate was approximately twice less than that of
RuOx/Ru(0001) under the same oxidizing conditions. This could be seen on Fig.3.17; it should
be taken into account that the reactivity of the clean Pt substrate must be subtracted since
the deposited film was grown on the one-side crystal. In analogy with a native oxide film a
spent deposited film maintained its ordering and a compositional stoichiometry according to
post characterization by LEED and AES.
The reason for such a low reactivity of RuOx/Pt(111) was studied in additional
experiments with pure oxygen treatment. Similarly to the native oxide films on Ru(0001)
freshly prepared and characterized RuOx films on Pt(111) were subjected to 20 mbar O2 at
450 K for 10 minutes in the HP cell. After this treatment they were thermally decomposed
and did not exhibit any O-species in low temperature interval (below 750 K) as depicted on
Fig.3.18 (compare with Fig.3.8 for native oxide films RuOx/Ru(0001)). Note that O-species
57
desorbing at ~ 1150 K after the main lattice O-peak at 1000 K belong to ‘as prepared’ oxide
film, i.e. not exposed to O2 in the HP cell. However, some new feature appeared: before the
lattice O-peak a broad moderate shoulder has developed at about 840-850 K. Its origin might
be interesting from structural point of view as an evidence of the surface reorganization.
Figure 3.17 Kinetics of CO oxidation over RuOx films grown on a one-side Pt(111) crystal and a double-
side Ru(0001) crystal. The reactivity of a clean one-side Pt(111) crystal is given for the reference. Thickness of
both oxide films is comparable.
Figure 3.18 TPD spectra of oxygen from decomposition in UHV of oxygen pretreated RuOx/Pt(111)
compared with RuOx/Ru(0001) (10 mbar O2 at 450 K, 20 min). Heating rate β=3 Ks-1
58
Observed results require additional morphological studies for the further
clarification. This finding apparently contradicts to previous, when presputtered native films
of ill-defined structure on Ru(0001) showed a pronounced enhancement of CO2 formation
with a following conclusion about the reaction to be structure insensitive.
Formation of an oxidic phase is not a prerequisite of high reactivity itself. However,
comparison of O2 TPD spectra shows that important is the desorption temperature of
oxygen species: similar to structure-reactivity dependence observed over FeO/Pt(111) [14-
15]. Indeed, good correlation between the increase of the film reactivity and the absolute
amount of weakly bound oxygen (WBO) as a function of film thickness allows to conclude
about higher availability of O atoms in RuO2(110) phase (growing preferentially on Ru(0001))
either upon film thickening and/or better ordering. Fig.3.19 represents data taken from
Figs.3.6 and 3.8 to make the statement more obvious: here a WBO-axis shows an integral
area under the low temperature O-peak (at 420 K) formed at elevated O2 pressures (Fig.3.8).
And since this peak develops only on oxide phase of (110) orientation due to its geometry,
oxide film thickness (and coverage) is not critical for the superior catalytic activity of
ruthenium oxide in the CO oxidation reaction.
Figure 3.19 Consistency between increase of the film reactivity and the absolute amount of weakly
bound oxygen (WBO) as a function of film thickness. The plots are offset from each other for clarity.
59
3.4 Summary
Thin ruthenium oxide films grown on Ru(0001) and O/Ru(0001) overlayers
were studied in the CO oxidation reaction at near-atmospheric pressures at
low temperatures (400 – 470 K).
It was shown that the reaction sets in only in the presence of the oxidic layer.
For the ultra-thin films (1-2 MLE), the surface exposed both the oxide and
O/Ru adsorbates structures. In addition, one-dimensional oxide structures
were observed, which were tentatively assigned to the intermediate state for
a crystalline oxide thin film that covers the whole surface at higher oxygen
coverage.
The films apparently maintained the structure and composition under highly
oxidizing reaction conditions.
The reaction rate slightly increases with increasing the nominal film thickness
and correlates with the total amount of weakly bound oxygen species formed
on the surface.
The disordered native oxide films on Ru(0001) showed even higher reactivity
than ordered counterparts. Therefore, it was proposed that surface ordering
of oxide phase does not play a key role in CO2 formation.
However, deposited RuOx films on Pt(111) possessing no (110) ordering and
weakly bound O-species have demonstrated much less reactivity towards CO
oxidation than RuOx films on Ru(0001) of comparable thickness under the
same oxidizing conditions.
Thus, obtained results suggest that the epitaxial growth of RuO2(110) phase is
a prerequisite for the Ru-based catalyst to be active and defines the reaction
mechanism. Possibly, the disordered RuOx films on Ru(0001) still maintained
some islands or domains of the active RuO2(110) phase, but they were too
small to be detected by LEED.
60
Chapter 4
CO oxidation over zinc oxide films on Pt(111)
4.1 Catalytic activity of ZnO-based materials and
growth of ZnO films
Besides large-scale use of ZnO as a semiconductor, colorant and technological filler,
zinc oxide is widely represented in heterogeneous catalysis due to industrial application in
methanol synthesis and water gas shift reaction (WGSR) [110, 111].
Preparation and characterization of polycrystalline powdered Cu/ZnO/Al2O3 mixtures
and their reactivity measurements under technologically relevant conditions are in very
detail described in works of R. Schlögl's and M. Muhler's groups [112, 113]. There the
authors came to the conclusion about the active involvement of metal-oxide interface in the
reaction mechanism when Cu particles distributed over an oxide support under reaction
conditions are partially covered with ZnOx species as a result of the SMSI effect.
The growth, electronic properties, interaction with a number of adsorbates and the
reactivity of Cu nanoparticles deposited on O-terminated ZnO(0001) (Zn- as well) single
crystals are reviewed in papers of C.T. Campbell with coworkers, e.g. [114] with references
herein. In parallel they studied the kinetics of forward and reverse WGSR on clean Cu(111)
and (110) to model high-area Cu/ZnO catalysts [115]. An attempt to grow an 'inverted'
ZnOx/Cu(111) model catalyst and employ it in methanol synthesis did not yield any product
formation, most likely because of the reaction conditions far from realistic [116].
Upon Zn-coverage dependent studies of CO2 hydrogenation over Zn-deposited low-
index Cu surfaces Nakamura et al. [117] concluded about structure sensitive Zn promotional
effect: in contrast to Zn/Cu(110) and Zn/Cu(100) systems Zn/Cu(111) showed high turnover
frequencies of methanol formation (18 atm, 523 K) at low Zn coverage (<0.2) due to Cu-Zn
surface alloy. At higher coverages (>0.2) Zn is selectively oxidized to ZnO leading to a drastic
61
decrease of the catalytic activity. In comparison with Cu(111) and other Cu planes the
activity of Zn/(111) was close to that of Cu/ZnO powder catalyst [118].
Involvement of metal-oxide interface in the reaction mechanism mentioned above
allows studying a model reaction, e.g. CO oxidation over a thin ZnO film grown on any
suitable (111) metal substrate as a good starting point.
Growth of the well-ordered ZnO films was previously reported using Ag(111) and
Pd(111) single crystal substrates. With the help of surface x-ray diffraction and scanning
tunneling microscopy (STM), Tusche et al. [119] observed that 2 ML-thick ZnO(0001) films
grown on Ag(111) are depolarized, i.e. Zn and O atoms are arranged in co-planar sheets like
in the hexagonal boron nitride (or graphite) structure. The transition to the bulk wurtzite
structure occurs in the 3–4 ML coverage range and is accompanied by considerable surface
roughening.
On Pd(111), Weirum et al. [120] observed the layer-by-layer growth mode for ZnO
films up to 5 ML, at least. Low energy electron diffraction (LEED) and STM results together
with density functional theory calculations suggested that the graphite-like structure is
thermodynamically the most stable phase over a large range of oxygen chemical potentials,
before it converges to the bulk-type wurtzite structure at a film thickness above 4 ML.
Small molecules can adsorb on individual clean ZnO phases depending on the surface
termination and polarity. This was intensively investigated by Wöll with coworkers and is
explicitly reviewed in [121]. Here they conclude referring to their own work [122] that
electrostatically unstable polar O-terminated ZnO phase is the most interesting from the
chemical point of view since it was theoretically predicted to be active in methanol synthesis
from syn-gas.
Although ZnO-based catalysts are primarily used in the methanol synthesis and the
water gas shift reaction, CO oxidation was previously studied as well on ZnO powders, single
crystals and films. In the work of Kobayashi with coworkers [123] with commercial ZnO
powder in addition to the Langmuir-Hinshelwood mechanism, i.e. between chemisorbed CO
and oxygen, and Eley-Rideal mechanism, with reaction between gaseous CO and adsorbed
62
oxygen, the simultaneous two reaction pathways mechanism was considered, in which two
different active species (neutral O and O-) were involved.
Esser with colleagues [124] reviewed earlier reports on the reactivity of
polycrystalline ZnO surfaces in CO oxidation and compared with that of ZnO(10-10) in own
experiments: they concluded that the results on polycrystalline samples are determined by
contributions from nonpolar ZnO(10-10) patches and that elementary steps of the reaction
can be explained on the basis of charge transfer.
Weiss and Folman [125] studied CO oxidation over ZnO films, fabricated by r.f.
sputtering on Ag foil, as a function of the film thickness. However, the thinnest films studied
were about 20 nm in thickness. Nonetheless, the authors found that the elementary
constants in the rate equations for CO oxidation at 600-650 K depend on the film thickness.
Good agreement between the experimental and the theoretical results, based on the
Schottky model together with the rigid band model, has been found.
Therefore, keeping in mind all the information above, here I have studied the
reactivity of ZnO films on Pt(111) in the CO oxidation reaction at low temperatures and near
atmospheric pressures in the continuation of previous works of our group on CO oxidation
over ultrathin oxide films.
To the best of my knowledge, the only available information about purposeful
growth of crystalline ZnO layers on Pt substrate refers to works about ZnO/Pt electrodes,
e.g. of Chugh with coworkers [126], where they reported the fabrication of
ZnO/Pt(111)/sapphire(0001) by PLD (pulsed layer deposition) and provided a HRTEM (high
resolution transmission electron microscopy) image of such a layered structure as well as
XRD (x-ray diffraction) proof of preferential formation of ZnO(0001).
Guided by ‘the electronic theory of catalysis’ and the concept of weakly-bond oxygen
both for FeO/Pt(111) and RuOx/Ru(0001) in this project I was driven by the idea of
establishing the relationship between the reactivity of ZnO films and their thickness.
63
4.2 Structure of materials and thin film preparation
Bulk ZnO can exist in the form of three polymorphs: zinc blend (sphalerite), wurtzite
and rock salt. Zn2+ and O2- ions are tetrahedrally coordinated to their counter ions in the first
two forms, while in the latter, which is stabilized at higher pressure, the local environment is
octahedral. The most thoroughly studied and thermodynamically stable under ambient
conditions phase is the wurtzite one. Polar (0001) or (000-1) as well as non polar (10-10) and
(11-20) planar surfaces can be prepared. Polarity of certain ZnO planes originates from the
ionic nature of Zn-O bond and, thus, an existing dipole moment and generated electric field.
Along the c-axis of the polar (0001) or (000-1) directions, ZnO exhibits repeating
Zn/O/… stacks of the hexagonal type. When the ZnO bulk is truncated normal to these
planes, the surface is Zn- or O-terminated, respectively, with three-fold coordinated atoms
(see Fig.4.1a). Such surface is electronically unstable, and surface charge is usually
compensated through surface reconstruction, facetting, electron redistribution or most
likely adsorption of foreign species from the residual atmosphere in the experimental set-up,
e.g. hydroxyls, yielding a lower surface energy. Stabilized ZnO(0001) and (000-1) basal faces
have a unit cell with a=3.25 Å in the real space giving a hexagonal (1×1) LEED pattern [96,
121, 127]. In contrast to ZnO single crystals with wurtzite structure ultrathin ZnO films grown
on Ag(111) [11] demonstrated planar graphene-like structure (Fig.4.1b).
Figure 4.1 Structures of ZnO bulk material and thin film: a - Stick-and-ball model of bulk ZnO of wurtzite
structure [96], b - graphene-like planar structure of 2 ML ZnO(0001) film on Ag(111): small and large balls refer
to Zn and O respectively [119].
a
b
64
Extended STM studies of geometric and electronic properties of low-index ZnO
surfaces of single crystals done by Diebold’s group [128, 129] revealed an essential
difference between the surface morphology of (0001)-Zn and (000-1)-O phases. The Zn-
terminated surfaces are characterized by a high roughness on the nanoscopic scale
exhibiting triangular islands and pits while oxygen-terminated surface presents flat
hexagonal terraces, wide and smooth with no added islands and holes.
Zinc was deposited on a double-side Pt(111) crystal by heating a Zn rod (1 mm in
diameter, 99.99%, Goodfellow) to ~ 500-520 K using W wire wrapped around. The Zn source
is shielded by a stainless steel cylinder having a small orifice (~ 5 mm in diameter) and placed
ca. 2 cm away from a crystal. The deposition flux was controlled via a Type K thermocouple
spot-welded to the edge of the Zn rod. The construction and the deposition recipe were
adapted from that of ZnO/Ag(111) of our group.
The zinc oxide films were prepared by Zn reactive deposition onto clean Pt(111) at 85
K or room temperature in 10-7 mbar O2 followed by oxidation at 600 K in 10-6 - 10-5 mbar O2
for 20 min. Low temperature Zn deposition favored more homogeneous Zn distribution over
Pt surface (to get a closed film in one step) and almost full Zn oxidation close to ZnO
stoichiometry prior to postoxidation.
Zn deposition in O2 ambient was carried out to prevent Zn dissolution in the Pt
substrate and any Zn-Pt alloy formation [130]. For most reactivity studies, oxide films were
grown on either side of the crystal. Such a method of Zn deposition from a metallic source in
an effusion cell further gives the oxide films of better quality in contrast to e-beam assisted
deposition from self-made stoichiometric ZnO target.
Postoxidation was necessary only for the film ordering, since required amount of
oxygen was already incorporated during the deposition step. Postoxidation temperature was
found experimentally to have an oxide film well-ordered and intact (see Figs.4.2 and 4.3).
The surface coverage was measured by CO titration of Pt sites (using TPD of CO that desorbs
from Pt(111) at 300 – 500 K) and AES and expressed in monolayers (ML) or monolayers
equivalent (MLE) (see below).
65
Figure 4.2 LEED patterns (at 60 eV) of ~1 MLE ZnO film on Pt(111), annealed in O2 at different
temperatures as indicated.
Figure 4.3 CO TPD profiles from clean Pt(111) and ZnO films on Pt(111) annealed at different
temperatures in O2 . Heating rate β=3 Ks-1
Figs.4.4 and 4.5 show LEED patterns and Auger spectra of the “as-prepared” films at
increasing thickness used for catalytic studies in CO oxidation. The diffraction spots of
ZnO(0001)-(1×1) are aligned with those of Pt(111) and show up together with the
surrounding hexagonal spots. The latter can be straightforwardly assigned to a coincidence
66
Pt(111)-(6×6) superstructure, similar to that observed on ZnO/Pd(111) [11], that arises due
to the mismatch between the lattice constants of ZnO(0001) and Pt(111) (3.25 and 2.78 Å,
respectively). Basically, 6 times the Pt(111) lattice coincide with 5 times the ZnO(0001)
lattice. Such geometrical superposition was already given in [131]. Above 2-3 ML, the films
only show a relatively diffuse pattern of ZnO(0001) with a high background intensity. The
ZnO-related spots were sometimes streaky in the “tangential” direction, thus indicating the
presence of domains slightly misaligned with respect to the metal support underneath.
Figure 4.4 LEED patterns (at 60 eV) of ZnO films grown on Pt(111) at the different film thickness in MLE as
indicated.
Figure 4.5 Auger spectra of ZnO films grown on Pt(111) at the different film thickness as indicated in MLE.
67
In submonolayer regime the surface coverage was measured by CO titration (see
Fig.4.6a). For more than a monolayer coverage, the nominal thickness was measured by
integrating the 32 amu (O2) desorption signal during film decomposition at T > 950 K (see
Fig.4.6b). Also the thickness was roughly monitored by AES via attenuation of the Pt related
signals with two restrictions: at submonolayer coverage sensitivity to Zn is too low for
precise estimation, at high coverage Zn and O signals stay constant with Pt signals almost
gone (see Fig.4.7).
Figure 4.6 ZnO/Pt(111) coverage calibration: a - CO TPD profiles from clean Pt(111) and ZnO films on Pt(111), b
- TPD spectra of molecular oxygen from decomposition in UHV of ‘as prepared’ ZnO films. Heating rate β=3 Ks-1
Figure 4.7 Auger signals of Pt and Zn as a function of the film thickness (in MLE). Characteristic breaking point
at ~ 1ML reflects layer-by-layer growth mode of ZnO films on Pt(111).
a
b
68
To address the morphology of the ZnO films additional structural characterization of
the similarly prepared oxide surfaces was performed in another UHV chamber equipped
with LEED/AES and high pressure scanning tunneling microscopy (HP-STM) by B.-H. Liu and
M.E. McBriarty, members of the 'Structure and reactivity' group, Department of Chemical
Physics, Fritz-Haber-Institute of Max Planck Society.
At sub-monolayer coverages, Pt(111) terraces were randomly covered by two-
dimensional islands (Fig. 4.8) [132]. The height measurements revealed that the islands are
predominantly of ~ 2 Å in apparent height, although few islands of ~ 4 Å in height were also
observed. These values correspond to one and two layers of graphite-like ZnO, respectively.
Henceforth, the amount of Zn to form a dense monolayer film will be referred to as one
monolayer equivalent (MLE).
Both mono- and bi-layer islands showed superstructure with a ~17 Å periodicity,
which is consistent with the Pt(111)-(6×6) superstructure observed by LEED. The films thicker
than 3 MLE showed no (6×6) superstructure in STM (also vanished in LEED), but triangular
pits and islands which are very similar to those observed on ZnO(0001) single crystal surfaces
in studies of Diebold’s group [128].
Figure 4.8 STM images of the “as prepared” ZnO films on Pt(111): 0.25 (a), 0.55 (b), 1.2 (c) and 4 MLE (d) [132].
The images are provided by B.-H. Liu and M.E. McBriarty.
Therefore, combined together the LEED, AES, STM and TPD characterization shows
that the “as prepared” films first form ZnO(0001) monolayer islands which coalesce until the
film covers the entire metal surface. The films then grow in a layer-by-layer mode. Upon
increasing the thickness above 4 MLE, the film surface is essentially the same as of a
ZnO(0001) single crystal, although the surface termination is not yet identified.
69
4.3 Results and discussion
4.3.1 Thickness dependence
The ZnO films, characterized by LEED, AES and CO TPD, were examined in the CO
oxidation reaction at near atmospheric pressures at 450 K. The reactivity measurements
were performed in excess of oxygen (CO/O2 = 1/5) in order to prevent any possible oxide
reduction and film dewetting, and also to compare with the previously studied FeO/Pt(111)
films, which exhibited film dewetting under oxygen-lean conditions [14], and RuOx/Ru(0001)
films.
Fig.4.9a displays typical reaction kinetics observed for the ZnO films showing no
deactivation in time. Fig.4.9b shows a CO2 production rate as a function of the nominal film
thickness measured prior to the reaction (overlapped reproduced data are not shown).
It is therefore clear, that the reaction exhibits a strong rate enhancement at sub-
monolayer coverage. Interestingly, the 1 MLE film shows some activity, at least, it is higher
than on pure Pt(111). As the thickness further increases, the rate gradually decreases such
that the multilayer, “thick” ZnO films become almost inert under the conditions studied. The
obtained trend is correct for ZnO films regardless the Pt(111) substrate temperature during
the deposition of Zn (85 K or room temperature ~ 300 K).
AES characterization of the spent catalysts did not reveal considerable changes in the
film stoichiometry, although some of the samples were lightly contaminated with carbon, in
particular at the sub-monolayer coverages.
LEED patterns showed diffuse (1×1) spots of Pt(111) and ZnO(0001), whereas the
(6×6) superstructure spots were vanished (see Fig. 4.10). Subsequent TPD measurements
showed a considerably higher CO uptake (during the first run after reaction is about twice
higher than before) as compared to that of measured prior to the reaction, thus indicating
the film dewetting and concomitant Pt surface opening under the reaction conditions (see
Fig.4.11).
70
Figure 4.9 CO oxidation over ZnO/Pt(111): a - Kinetic curves for the pure Pt(111) and ZnO(0001) films. Reaction
conditions: 10 mbar CO + 50 mbar O2, He balance to 1 bar; 450 K, heating rate 1 Ks-1 (to reach 450 K). b - CO2
production rate as a function of the nominal film thickness (in MLE).
Broadened and shifted to lower temperatures CO TPD profiles from postreacted films
during sequential TPD runs after the first one are very similar to that reported for bimetallic
surfaces of Zn/Pt(111) [130]. This could be ascribed to ZnO-Pt(111) alloying when heating
the spent sample from 85 K to 600 K in UHV: upon burning some carbonaceous deposits
formed under reaction conditions consume oxygen from the oxide film.
a
b
71
Figure 4.10 LEED patterns (60 eV) of the ZnO films on Pt(111) before (upper panel) and after (down panel) the
reaction. The film thickness increases from left to right: 0.7, 1 and 2 MLE.
Figure 4.11 CO TPD for as prepared and post reacted surface of a ZnO film, θ=0.7. Heating rate β=3 Ks-1
To examine whether the reaction is accompanied by Zn-Pt alloy formation, we
studied the Zn-Pt alloy. For this, ~ 1 ML of Zn was deposited on clean Pt(111) surface with
consequent annealing in UHV up to 600 K which resulted in the ordered (2×2) surface
structure (see Fig.4.12). As prepared sample demonstrated well-documented CO TPD profile
[130]. Explanation of such a well-pronounced shift of CO desorption peak from a freshly
72
prepared Zn-Pt(111) alloyed surface was given in [130]: the electronic perturbations induced
by Zn on Pt reduce its CO-chemisorption ability by weakening the strength of the Pt(5d)–
CO(2π*) bonding interactions.
Figure 4.12 LEED patterns (60 eV) of an alloyed Zn-Pt(111) surface before (a) and after (b) annealing to 600 K
The reactivity of initially alloyed surface was comparable with that of some
submonolayer ZnO films (see Fig.4.9a). And even post characterization by LEED did not show
any ordered pattern, AES definitely revealed ZnO oxide formation (see Fig.4.13) due to high
chemical potential of oxygen under reaction conditions. This conclusion is also supported by
the first CO TPD after reaction, which is similar to that of preformed oxide films (see Fig.4.14,
compare with Fig.4.11).
Figure 4.13 Auger spectra of an alloyed Zn-Pt(111) surface before and after reaction.
a
b
73
Figure 4.14 CO TPD for as prepared and post reacted surface of an alloy Zn-Pt(111). Heating rate β=3 Ks-1
A temperature programmed reaction (TPR) was carried out (see Fig.4.15) over 1 MLE
ZnO film with the nominal thickness ~ 1 ML to detect the point, when the reaction ignites:
based on experiment an estimated value lies above 420 K. It is noteworthy that carbon
contamination on the surface of the spent oxide film after TPR was below AES detection
limit. This could be explained by constant increase of oxygen partial pressure during the
reaction when CO gradually converts to CO2, and, therefore, carbon deposits forming in situ
are burnt.
Figure 4.15 CO2 production measured during TPR over 1 MLE ZnO film on Pt(111). Reaction conditions: 10 mbar
CO + 50 mbar O2, He balance to 1 bar; T=300-470 K, heating rate 0.02 Ks-1 (during the reaction). Here one-side
polished Pt(111) crystal was used as a substrate.
74
Also the LEED pattern certified maintenance of the (6×6) superstructure (see
Fig.4.16a and b) – at least in (0;0) in contrast to the case of a ‘normal’ catalytic run when this
superstructure was vanished (see Fig.4.10). This can be a sign of the fact that thin ZnO films
are able to restore their stoichiometry under oxidative conditions.
Figure 4.16 LEED patterns (60 eV) of 1 MLE ZnO film on Pt(111) before (a) and after (b) TPR.
4.3.2 Active phase formation and structural stability of ZnO films
In analogy with earlier works of our group on CO oxidation over ultrathin oxide films
[14, 100] it was first proposed that the reactivity of ZnO films on Pt(111) possibly could be
related to the new oxygen species formation under reaction conditions in excess of oxygen.
To examine this, freshly prepared ZnO films (0.9, 1.1, 2.2, 5 MLE thick) were subjected to
similar O2 treatment: 20 mbar of pure O2, 450 K, 10 min.
Thermal decomposition of the films evidenced only the presence of the lattice
oxygen peak at ca. 1030-1100 K. Closed ZnO films did not possess any additional oxygen
species. Partially covered films demonstrated a small extra-peak at lower temperatures, ca.
740 K. However, such a peak can be easily attributed to O2 desorption from opened patches
of the Pt(111) surface: this is well-documented in the work of Steininger with colleagues
[133].
The morphology of the catalysts was studied by B.-H. Liu with STM ex situ [132]. For
this, the well-defined films were exposed to the same reaction conditions (10 mbar CO, 50
mbar O2, He balance, 450 K) in a high-pressure cell for 10 min. The sample was evacuated
and transferred into analytical chamber via a gate valve and immediately scanned with STM
at room temperature. The STM images of the “reacted” samples are displayed in Fig.4.17.
a
b
75
For the 0.25 MLE film, irregularly shaped islands of ~10 nm in lateral size are observed
(Fig.17a). The film coverage is reduced to ~ 0.18 ML (which is most likely overestimated due
to the well-known tip convolution effect for small objects). No monolayer islands, which
dominated the surface of the “as prepared” film (see Fig. 4.8a), are observed anymore.
Instead, the particles are all about 4 Å in height, thus indicating the formation of two-
layers thick (henceforth “bilayer”) ZnO(0001) islands. Basically, similar behavior is observed
for the 0.55 MLE film (Fig.4.17b): The surface coverage decreases to ~ 0.32 ML, and the
Pt(111) support is covered with small ~ 4 Å-high islands and their aggregates. Finally, the 1.2
MLE film dewets as well, thus forming a 0.75 ML covered “bilayer” film (Fig. 4.17c).
For the all films inspected, only small amounts of ZnO are found in the third layer.
The long range ordering via the Moire (6×6) superstructure seen on the “as prepared” films
is hardly visible on the “reacted” films, in full agreement with the LEED data (Fig. 4.10).
Figure 4.17 STM images of the “reacted” ZnO films at 0.25 MLE (a), 0.55 MLE (b) and 1.2 MLE (c) film
thicknesses [132]. The images are provided by B.-H. Liu and M.E. McBriarty.
Observed structural change of 'postreacted' thin film surfaces correlates well with
increased CO desorption from that samples as revealed by CO TPD. On Fig.4.18 (is presented
with another focus than Fig.4.11) one can clearly see that ZnO dewets the Pt substrate under
reaction conditions. The degree of dewetting is hard to quantify precisely because of the
presence of carbonaceous deposits on the spent catalysts also contributes to CO signal by
reaction with oxygen available in the system as it could be seen from the baseline rising up
at T > 500 K.
a
b
c
76
Figure 4.18 CO TPD spectra of the 0.7 ML film before and after the CO oxidation reaction: film dewetting. The
spectrum of the clean Pt(111) is given as a reference. β=3 Ks-1
Therefore, the STM results show that the CO oxidation reaction on ZnO/Pt is
accompanied by substantial structural transformations such that monolayer islands (and a
dense monolayer film) transform into the bilayer islands and their aggregates. Therefore, Pt
surface opening observed by CO uptake measurements of the spent catalyst reflects, in fact,
a higher stability of the bilayer structure at elevated pressures.
The exact mechanism and the driving force for the observed structural changes are
hard to envision on the basis of the existing results. Note, however, that bilayer ZnO islands
were also developed on Pd(111) at increasing oxygen pressures used for oxidation at 550 K
[120]. Interestingly, bilayer islands were observed for CeO2-x(111) deposited onto Rh(111)
[134] and Ru(0001) [135].
The rate enhancement on the metal surfaces partially covered by oxide is commonly
rationalized in terms of the reaction occurring at the oxide/metal boundary. Indeed, the
maximum in activity basically follows the perimeter length of the oxide/metal interface
which goes to zero for a dense oxide film. The effect may result either from oxygen spillover
from oxide to Pt or from the formation of highly active sites at the metal/oxide boundary.
Since ZnO oxide is non-reducible oxide, the spillover mechanism seems to be not operative.
Previously, Hardacre et al. [136] found maximum in low temperature CO oxidation on
CeOx/Pt(111) in the sub-monolayer regime. The coverage was measured by CO titration of
77
bare Pt(111). Reactivity was measured in a high-pressure cell used as a batch reactor filled
with ~ 10 Torr of stoichiometric (2:1) CO + O2 mixture at crystal temperature of 320 – 430 K.
The authors found that the reaction rate increased between zero and 0.5 ML, after which it
dropped almost to zero at 0.8-1.3 ML, but again increased steeply to a value which is greater
than that observed over the clean Pt(111) surface.
It was suggested that for the coverages below 1 ML the reactivity is associated with
the metal/oxide interface. The high reactivity of the fully encapsulated Pt(111) (as judged by
CO TPD before the reaction) was rationalized in terms of a proposal made by Frost [137],
according to which electron transfer from a metal phase to an oxide phase reduces the
enthalpy for oxygen vacancy formation in the oxide. However, the authors argued
themselves that there was no absolute evidence of total encapsulation of the Pt crystal. In
particular, the sample morphology after the reaction was not defined in this work. The
initially dense oxide film could dewet, ultimately forming CeOx particles on Pt(111).
Nonetheless, the rate vs thickness plot observed for CeO2-x(111)/Pt(111) system is
very different from that shown in Fig.4.9b for ZnO(0001)/Pt(111). While “thick” ceria films (~
10 MLE) were much more active than Pt(111), the reactivity of ZnO(0001)/Pt(111) dies away
with increasing the film thickness. The difference may be related to the fact that ceria, well-
known for the oxygen storage-release properties, readily provides weakly bound oxygen to
react with CO, whereas ZnO as non-reducible oxide does not.
Recently Sun et al. [138] performed density functional theory studies of the 3d
transition-metal oxide (TMO) nano-islands on Pt(111) in low-temperature CO oxidation in
order to identify the active sites at the TMO/Pt boundaries. A Pt–cation ensemble was
proposed, where coordinatively unsaturated TMO cations exposed at the edges of oxide
nanoislands are highly active for O2 adsorption and dissociation, and less-reactive Pt binds
modestly with dissociated O responsible for the facile CO oxidation. But, in all these
calculations the TMO islands were only one monolayer in thickness, which is not the case
under technically relevant conditions as shown here for ZnO(0001) and previously for
FeO(111) on Pt(111) [14]. Therefore, our results provide more adequate structural models
for TMO/metal “inversed” catalysts.
78
The so called ‘inverse’ catalysts in the CO oxidation reaction, when reactants interact
on the oxide-metal interface, are already well-known in the literature, e.g. for such cases like
VOx/Pd(111), VOx/Rh(111), CeOx/Pt(111), SnOx/Pt(111) and CeOx/Cu(111) respectively [139-
143]. However, the first four reported studies utilize low pressure reaction conditions. In the
work with CeOx/Cu(111) [143] Yang et al. performed the reaction at elevated pressures (20
Torr CO, 10 Torr O2, 475-575 K), but formation of copper oxide phases obscured establishing
the active phase.
The idea of the active oxide/metal interface is a wide explanation and can include, at
least, three potential pathways. These suggestions are roughly depicted on Fig.4.19 and
marked respectively as 1, 2 and 3. The most reasonable one can involve CO adsorption on Pt
sites and reaction with oxygen adsorbed on oxide film or coming from its lattice to produce
CO2 (1). Another scenario can be opposite to the first one: CO adsorbs on some defect sites
of ZnO (e.g. Zn or O vacancies) and further interacts with atomic oxygen dissociated on Pt
(2). Even if the electron transfer seems to be not the case for ZnO/Pt(111), as it was proven
experimentally and theoretically for FeO/Pt(111), the third way can be true when Zn-Pt pair
(not an alloy) as a local junction creates electronically modified area, and both reactants can
easier adsorb at the oxide/metal boundary and combine with each other (3).
Figure 4.19 Schematic representation of possible pathways of the reaction mechanism of CO oxidation
over ZnO/Pt(111) bilayer islands.
Thermal stability of ZnO films was studied by TPD. Experiments conducted with the
films prepared on two different Pt(111) substrates (one- and two-side polished single
crystals) have shown the same trend: the leading edge of the molecular oxygen peak for
films thicker than 2 MLE is shifted to higher temperatures (see Fig.4.20). This could serve as
79
an evidence of substantial difference between thin and relatively thick films, being a sign of
strained geometry: the stress, originated from the lattice mismatch between Pt(111) and
ZnO(0001) and accumulated in thin films, releases upon thickness increase. A transition
thickness of ca. 2 MLE differentiates ZnO ultrathin films with particular properties from
already bulk-like thicker films. This issue needs further investigations.
Figure 4.20 TPD spectra of molecular oxygen from decomposition in UHV of ‘as prepared’ ZnO films grown on
two different Pt(111) crystals. Heating rate β=3 Ks-1
Also in this respect a question of surface termination might be interesting, even
observed for ZnO single crystals. Previously cited STM studies of Diebold’s group certified
the height difference between Zn- and O-terminated ZnO(0001) surface [128]: ~ 2.7 Å single
step for (0001)-Zn against ~ 5.3 Å for double layer step of (000-1)-O phase.
Moreover, Lauritsen et al. in their combined study with high-resolution scanning
probe microscopy (SPM), X-ray photoelectron spectroscopy (XPS), near edge X-ray
absorption fine structure (NEXAFS) spectroscopy experiments, and density functional theory
(DFT) calculations explain higher probability of existence of (000-1)-O surface by a higher
bonding flexibility of the exposed 3-fold coordinated surface Zn atoms as compared to O
atoms [144].
80
4.4 Summary
Well-ordered ultrathin ZnO(0001) films on Pt(111) are shown to grow in a
layer-by-layer mode. The films are examined in the CO oxidation reaction as a
function of the film thickness. At low temperatures (~450 K) and near-
atmospheric pressures, CO2 production is much higher on the partially
covered films than on the dense ZnO(0001) films and bare Pt(111).
Under reaction conditions, both monolayer islands and a monolayer film
transform into bilayer islands, which dominate the surface of the active
catalysts. The results provide more adequate structural model than previously
suggested for elucidating the reaction mechanism on the active oxide/metal
boundary at technically relevant conditions. Stability of the ZnO bilayer
structure at elevated pressures shows that the active phase differs from the
pristine one stable and observed under UHV conditions. Therefore, the
present work demonstrates, how the ZnO surface may respond dynamically
to changes in the surrounding ambient atmosphere at elevated pressures.
81
Chapter 5
CO oxidation over manganese oxide films on Pt(111)
5.1 Catalytic activity of MnOx-based materials and
growth of MnOx films
Besides industrial applications of various bulk manganese oxides, e.g. related to
energy storage [145-147], they as high area porous crystalline materials have an own niche
in heterogeneous catalysis. Manganese oxide was reported to promote CH4 dry reforming by
CO2 to CO and H2 (syngas) through mobility facilitation of bulk lattice oxygen [148]. Stabilized
by alkaline additives MnOx were shown to be active in oxidative coupling of methane (OCM),
where strongly bound (lattice) oxygen is responsible for selective methane coupling, while
weakly bound oxygen - for conversion of CH4 to CO2 (exhaustive oxidation) [149, 150].
Manganese oxides are known to be active catalysts in ethylene [151] and alcohols
[152] oxidation; NO [153], H2O2 [154] and O3 [155] decomposition; selective catalytic
oxidation of NH3 [156], oxidative dehydrogenation ethylbenzene to styrene [157], selective
catalytic reduction of NO with NH3 [158], CO [159] and C2H4 [160] hydrogenation. The
catalytic activity has been attributed to the capability of Mn to form several oxides and to
store and provide oxygen selectively from its lattice. Because of its labile oxidation state, Mn
is capable of playing the role of either a reducing agent that is oxidized (Mn2+ - ē → Mn3+ - ē
→ Mn4+) or an oxidizing agent that is reduced (Mn4+ + ē → Mn3+ + ē → Mn2+).
MnOx-based catalysts are utilized for industrial wastewater treatment through
catalytic wet air oxidation to degrade harmful organic pollutants [161]. Another environment
related use of MnOx compounds refers to oxidation of volatile organic compounds (VOCs)
[162]. In that work authors proposed that lattice oxygen was involved in the VOCs oxidation,
suggesting that the reaction could proceed via the Mars-van Krevelen mechanism.
Additionally, the conversion was found to be influenced by the type of VOCs and catalysts
provided by the shape-selectivity of MnOx porous materials, octahedral molecular sieves
(OMS) [163, 164]. Manganese oxides nanoparticles can biomimetically catalyze water
oxidation in artificial photosynthesis to produce molecular oxygen O2 [165].
82
Catalytic CO oxidation over various kinds of individual and mixed manganese oxide
phases is in general overviewed by Royer and Duprez [166]. In detailed experimental studies
of Kruse’s group with MnOx nanocrystals [167] the authors conclude that key requirement
for catalytic activity is nanosize of MnOx particles (as proven by transmission electron
microscopy (TEM)) and, thus, high specific surface area of the catalysts. Based on CO
oxidation in absence of O2 it was proposed that due to excess of oxygen in surface region the
Mars-van Krevelen type reaction mechanism is most likely in operative.
This is in agreement with the work of Baltanas et al. with highly dispersed overlayers
of MnOx [168]. They found that the reaction has the first order for CO and O2. The amount of
active oxygen on the surface is directly proportional to the catalytic activity, and is less than
10% of the total surface oxygen. The authors highlight that this active oxygen does not
correspond to a labile moiety but is a part of the structural oxygen of the surface.
To differentiate several forms of oxygen bonds on the surface of oxide solids
including manganese oxides, oxygen isotope exchange was utilized using O18-labeled O2, CO
and CO2 molecules, as it is reviewed by Novakova [169]. Recent studies of Fishman et al.
[170] with MnOx ‘nanopowders’ supported known regularities on the new level.
Catalytic properties of MnOx-based materials mentioned above refer to studies
where the catalysts were prepared by ‘wet chemical’ methods and tested either at
atmospheric pressures or higher and over the wide temperature range (mostly elevated).
Single crystalline MnOx surfaces or metal-supported thin films have never been reported to
be an object of studying any catalytic reaction. However, there is plenty of literature focused
on growth and structure of MnOx thin films of different thickness and stoichiometry.
Manganese oxide thin films were reported epitaxially to grow on a number of well-
known substrates in surface science, namely on Ag [171], Pd [172], Pt [173] and Rh [174].
Upon multitechnique research (with LEED, X-ray photoemission spectroscopy (XPS), STM, X-
ray photoelectron diffraction (XPD) and X-ray absorption spectroscopy (XAS)) on
MnO(001)/Ag(001) it was found that MnO films few nm thick are strained and tetragonally
distorted, while the thicker films evidently adopt bulk structure and bulk lattice constants
[175-177]. However, MnO films can intermix with Ag(001), forming embedded layers [178].
83
A variety on MnOx phases (in total nine) of different orientation and stoichiometry
(e.g. MnO - (001), (100), (111), Mn3O4 - (001)) was revealed on Pd(100) as a function of the
chemical potential of oxygen and thickness in works of Netzer’s group [172, 179-180], where
the data were collected from LEED, STM, XPS, XAS, high-resolution electron energy loss
spectroscopy (HREELS) and atomic force microscopy (AFM) methods.
In studies of Widdra’s group [173, 181] vibrational and electronic properties of
manganese oxide films grown on Pt(111) were presented to be thickness dependent, as
proved by HREELS technique: new excitations are observed below 4 ML, while films adopt
the bulk-like structure comparable with that of single crystals already from the thickness ca.
4 ML. Additional LEED and STM characterization and TPD measurements demonstrated
existence of different phases as a function of O2 partial pressure and the substrate
temperature during the film preparation. The film coverage was calibrated by CO TPD.
The hexagonal structure of MnO monolayer on Rh(100) was found and explained by
Nakamura’s group [174], based on LEED and STM studies. Such model system was proposed
potentially to have a meaning for CO hydrogenation reaction. Recently, by Zhang et al. [182]
MnOx thin films were shown to grow on Rh(111) as an O-Mn-O like trilayer, above that
coverage it continued to develop as Mn3O4 structure, as detected by combination of LEED,
AES, XPS, HREELS and low energy ion scattering spectroscopy (LEIS) techniques.
It is worth to underline, that, in spite of availability of numerous studies regarding
structural properties of MnOx thin films, the information about small molecules adsorption
in UHV on ordered individual MnOx surfaces is still missing.
The idea about highly reactive surface oxygen species was already clearly introduced
for the bulk and supported manganese oxides. Therefore, for me in the current work it was
reasonable to perform catalytic studies of a model reaction, e.g. CO oxidation over ultrathin
MnOx films grown on a suitable substrate like Pt(111) in order to detect and prove the
formation of new oxygen species under well-defined conditions, identify their nature and
establish thickness-reactivity dependence if any. Moreover, the comparison of the reactivity
of MnOx/Pt(111) model catalyst at near-atmospheric pressures and low temperatures with
that of FeO/Pt(111), RuOx/Ru(0001) and ZnO/Pt(111) sounds interesting.
84
5.2 Structure of materials and thin film preparation
As a 3d-transition metal manganese with an electronic configuration [Ar] 4s2 3d5 can
form a large variety of oxides with the oxidation state between +2 and +7. Most commonly
known of them are crystalline polymorphs of MnO, Mn3O4, Mn2O3 and MnO2 stoichiometry,
which basic data are compiled in Tabl. 5.1 [145, 183-185] (their literature nomenclature is
often contradictory). Corresponding structures of most stable allotrope modifications are
shown below in Fig.5.1. Here I overview materials consisting only of Mn and O ions, since
detailed description of geometry of substituted forms, hydrates and minerals is beyond the
scope of the present work. Temperature dependent magnetic properties of manganese
oxides, which are intimately connected to materials geometry, are not overviewed as well.
From my point of view such extended structural information about bulk manganese oxides
as compared to corresponding sections in previous chapters is reasonable since MnOx show
much higher complexity in this respect than ruthenium and zinc oxides.
Figure 5.1 Schematic representation of bulk manganese oxides [145]. (a) Rock salt; (b) spinel (Mn3O4);
(c) bixbyite (Mn2O3); (d) pyrolusite β-MnO2 (rutile -type); (e) ramsdellite; (f) phyllomanganate
85
Table 5.1 Crystal structure of bulk manganese oxides
Formula
Trivial name
Symmetry
Structure type
Unit cell
Cleavage
a
b
c
β
MnO
manganosite
cubic
Rock salt (at room temperature) - face centered cubic (FCC) lattice
with a 6 : 6 octahedral coordination.
4.44
100, 010,
001, 111
rhombohedral (4.2 K)
4.43
Mn3O4
α-, hausmannite
tetragonal distorted
Spinel - metal cations occupy 1/8 of the tetrahedral sites and 1/2 of the
octahedral sites and there are 32 oxygen anions in the FCC unit cell.
5.76
9.47
001, 100,
110, 010
β-
cubic (≥ 11700C)
spinel
8.55
Mn2O3
α-, kurnakite
orthorhombic
distorted
9.41
9.42
9.40
β-, bixbyite
cubic (≥ 302 K)
Body-centered cubic (BCC) unit cell with 16 formula units per unit cell. Fe
ions
9.41
110, 111
γ-
spinel
MnO2
α-, hollandite
monoclinic
Cross-linking of double or triple chains of the [MnO6] octahedra,
resulting in (2×2) opened tunnels within the lattice. Ba ions
10.03
2.88
9.73
91.03
α-, cryptomelane
monoclinic
-//-. K ions
9.96
2.87
9.71
90.95
β-, pyrolusite
tetragonal
Rutile structure with an infinite single chain of [MnO6] octahedra running
along the [001] direction and sharing opposite edges; each chain is
corner-linked with four similar chains, forming (1×1) square cavities.
4.40
2.87
110, 100,
101
R-, ramsdellite
orthorhombic
Closely related to rutile except that the single chains of edge-sharing
octahedra are replaced by double chains, resulting in (1×2) channels.
9.27
2.87
4.53
γ-, nsutite
hexagonal
An irregular intergrowth of layers of pyrolusite and ramsdellite.
9.65
4.43
δ-, vernadite
preuso(tetragonal)
Layered structure, containing infinite two-dimensional sheets of edge-
shared [MnO6] octahedra. nH2O
9.87
2.84
ε-, akhtenskite
hexagonal
Defective NiAs hexagonal close packing of anions, with Mn4+ statistically
distributed over half the available octahedral interstices
2.84
4.6
λ-
cubic
Spinel with channels
8.04
η-
Different from γ -MnO2 only in crystallite size and the concentration of
microdomains of pyrolusite within the ramsdellite matrix
86
There are not so many known experimental studies of the surface properties of
manganese oxides single crystals due to charging problems. However, Langell et al. reported
selective oxidation/reduction of Mn(100) substrate and epitaxial growth of Mn2O3 and Mn3O4
like surfaces under UHV conditions as supported by HREELS and XPS characterization [186]. The
authors conclude that the ready formation and intermixture of these species imply that the
redox properties of manganese oxide surfaces are important in their surface chemical reactivity
and gas adsorption characteristics.
Later, unreconstructed and clean Mn(001)-(1×1) surface was successfully prepared by
cleavage in N2 ambience by Okazawa and Kido [187] and analyzed with reflection high energy
electron diffraction (RHEED) and high-resolution medium energy ion scattering (MEIS)
techniques. Long annealing in O2 at too high temperatures and pressures led to the surface
amorphization. Annealing at lower O2 pressures generated oxygen vacancies at the surface.
Formation of Mn3O4(001) phase on MnO(001) surface was shown in experimental-
theoretical study by Bayer et al. [179] (methods are indicated in Sec.5.1). Authors computed
low strain energy and proposed realistic models for the interface, even with 9% lattice
mismatch between two phases on the interface.
Gorbenko et al. [188] grew thick ordered films of Mn3O4 (100 nm) on MgO(001) and
SrTiO3(001) single crystals and characterized them by X-ray diffraction (XRD), Raman
spectrometry (RS) and high-resolution transmission electron microscopy (HRTEM). On MgO
(001) only domains of (001)[100]Mn3O4//(001)[100]MgO symmetry were observed, while on
SrTiO3(001) different orientations were detected: (001)[110]Mn3O4//(001)[100]SrTiO3,
(001)[110]Mn3O4//(001)[100]SrTiO3 and (101)[010]Mn3O4//(001)[110]SrTiO3.
Natural pyrolusite and synthetic MnO2 material surfaces were both shown to crystallize
in [110] direction, meanwhile exposing other planes as well according to HRTEM images of
Yamada et al [189] and Zheng et al. [190] respectively.
87
Bulk-terminated polar MnO(111) phase first was calculated with DFT by Franchini et al.
[191]. They reported unreconstructed O-terminated (1×1)-MnO(111) surface to be stable at
high oxygen pressures. Previously MnO(111) thick films (200-3500 Å) were epitaxially grown on
sapphire(001) and MgO(111) by a pulsed layer deposition (PLD) by Neubeck et al. [192]. Films
orientation was proven by XRD, grazing incidence diffraction (GID), X-ray reflectometry and
Rutherford Back Scattering (RBS) methods.
Under some preparation conditions an oxygen-terminated monolayer film O-MnO(111)
was suggested to exist on Pt(111) substrate in a form of a O-Mn-O trilayer (see Fig.5.2 for
clarity) with MnO2 stoichiometry by Sachert (Widdra’s group) [193]. Top layer of oxygen was
found to be weakly bound as compared to MnO(19×1)-Pt(111).
Figure 5.2 Structural model of an oxidized MnO monolayer film on Pt(111). Large and small balls refer to Mn and O
atoms correspondingly [193].
In this work thin manganese oxide films of different structure and stoichiometry were
prepared using conditions adapted from [173, 181, 193]. Manganese was evaporated at room
temperature on a biased Pt(111) crystal by e-beam assisted (reactive) physical vapor deposition
(PVD) in either UHV or O2 ambient (~ 5·10-8 mbar) from a metallic Mn rod (2 mm in diameter,
99.5%, Goodfellow). The deposition flux was controlled via a quartz microbalance (QMB).
'O-poor' MnOx phases were obtained by postoxidation of Mn deposits at 700-1000 K for
several minutes in 10-6 mbar O2. Structures with higher oxygen content were produced by
oxidation of Mn deposits in 10-4 mbar O2 at 700 K for 5 min. 'O-rich' MnOx phase was prepared
by oxygen treatment of any ‘O-poor’ phase in 20 mbar O2 at 450 K for 10 min in the HP cell.
88
Monolayer films deposited in UHV and annealed in O2 below 850 K demonstrated a
(19×1) LEED pattern and Mn:O=1:1 stoichiometry in AES (that corresponds to the peak ratio of
1.67) as shown on Figs.5.3a and 5.4. The (19×1) LEED pattern is a result of superposition of
three 120o-rotational domains with (2.37×1) periodicity as proposed in [173]. Therefore, based
on such characterization fingerprints, films were attributed to MnO(100) monolayer uniaxially
reconstructed on Pt(111).
O2-deposited films annealed in O2 at T ≥ 850 K displayed a LEED pattern (see Fig.5.3b)
not yet identidied due to its complexity. Elemental composition yielded Mn:O ratio ranging
between 1.2 and 1.4 (see Fig.5.4, overlaps with MnO). A phase with these characteristics was
assigned [193] to Mn3O4. It can convert to MnO(100) structure upon heating in UHV to 1190 K
or obtained by MnO reoxidation in O2 ambient at T>850 K.
O-rich phase formed from either MnO or Mn3O4 under mbar O2 pressure at 450 K
exhibited a diffused (1×1) LEED pattern with high background intensity. This could be explained
by the oxide surface amorphization, while Pt spots at 1 ML film thickness are still visible (see
Fig.5.3c). Or domains of an ordered structure are too small and have multiple orientation.
Meantime, the chemical potential of oxygen was high enough to result in double increase of its
amount and, therefore, in MnO2 stoichiometry as judged by AES (see Fig.5.4). Direct deposited
Mn oxidation in 10-4 mbar O2 at 700 K led to a structure with a similar pattern as in HP cell,
however, slight onset of a weak Moiré superstructure could be distinguished (see Fig.5.3d). If
latter would be fully developed, it could be assigned to O-Mn-O trilayer of O-MnO(111)
possessing a well-pronounced (1×1) Moiré pattern [49].
Figure 5.3 LEED patterns of 1 ML MnOx films on Pt(111) at 70 eV: a - MnO, b - Mn3O4 and c(RT), d(700 K) - MnO2.
a b c d
89
Figure 5.4 AES of 1 ML MnOx/Pt(111) subsequently oxidized through MnO, Mn3O4 and MnO2 steps.
The surface coverage was measured by CO titration of Pt sites using TPD of CO that
desorbs from Pt(111) at 300 – 500 K.
5.3 Results and discussion
5.3.1 Active phase formation and special requirements for its existence.
Catalytic reactivity of bulk and supported MnOx materials in various chemical processes
overviewed in Sec.5.1 as well as structural studies of O-MnO(111) film on Pt(111) [193] already
documented existence of other forms of oxygen different from oxide lattice. Therefore, prior to
start studying the reactivity, in a separate series of experiments the ‘as prepared’ 1 ML films
were exposed to O2 in HP cell and then thermally decomposed in UHV. A TPD profile depicted
on Fig.5.5 clearly indicates three peaks of molecular oxygen desorption - 950, 1060 and 1230 K
respectively. Peak positions allow to conclude that chemical composition is virtually similar to
that of the O-Mn-O trilayer [193].
90
Figure 5.5 TPD profile of MnOx/Pt(111) pretreated with O2 (20 mbar, 450 K, 10 min), β=3 Ks-1.
However, peaks integrated area ratio confirms the coexistence of MnO2 and Mn3O4
rather than pure MnO2 in contrast to [193] where the presence of Mn3O4 in the TPD spectrum
is almost negligible. It means that during decomposition upon losing oxygen MnO2 is followed
by Mn3O4 and finally MnO.
Initially the reactivity measurements for CO oxidation were performed under the same
conditions used for other oxide films in the present work (RuOx/Ru(0001) and RuOx/Pt(111),
ZnO/Pt(111)) and earlier studies of our group (FeO/Pt(111)): namely in a mixture 10 mbar
CO+50 mbar O2, balanced by He up to 1 bar, at 450 K. On Fig.5.6 a kinetic curve of CO2
production over a monolayer MnOx film is compared with that of a monolayer FeO film and
bare Pt(111) surface. The reactivity of MnOx/Pt(111) is lower than over the FeO film and just
slightly above that of the Pt(111) substrate. After a short working period, CO oxidation slows
down and stops.
Characterization of the postreacted oxide surface has shown that the film underwent
severe dewetting as well as O-depletion accompanied by C deposits. This was concluded from a
characteristic shape of the CO TPD profile (see Fig.5.7a) and a Pt(111)-c(4×2)CO LEED pattern
91
(see Fig.5.8a), both comparable with that of postreacted Pt(111) in a blank experiment.
Subsequent heating to 600 K in the CO TPD run led to wetting of the most part of the Pt surface
by MnOx film with the development of the (19×1) LEED pattern (see Fig.5.8b). The film
decomposition revealed no additional O2-peaks besides the lattice one at ~ 1200 K (see
Fig.5.7b).
Figure 5.6 Kinetics of CO oxidation over 1 ML MnOx film on Pt(111) at CO:O2=1:5. Reaction conditions:
10mbar CO+50 mbar O2, balanced by He to 1 bar; 450 K, heating rate 1 Ks-1 (to reach 450 K). Kinetic curves for 1
ML FeO film on Pt(111) and clean Pt(111) are given as a reference.
Figure 5.7 CO TPD (a) of 1 ML MnOx/Pt(111) and its decomposition (b) after reaction with CO/O2=1:5, β=3 Ks-1.
b
a
92
Figure 5.8 LEED patterns (70 eV) of a spent 1 ML MnOx/Pt(111) after the reaction CO/O2=1:5 (a) and CO TPD (b).
In attempts to prevent film dewetting, in the next experiments we increased CO/O2 ratio
from 1:5 to 1:10, keeping the total pressure of reactants to 60 mbar. Upon increase of CO:O2 to
1:10 in the reaction gas mixture, the film catalyzes CO oxidation with a much higher and
constant reaction rate (see Fig.5.9). CO titration experiments confirmed that the film does not
dewet (see Fig.5.10a). Additionally, during thermal decomposition in UHV of postreacted oxide
surface the profile of O2 spectrum demonstrates O-states similar to that observed for the
surface pretreated in pure oxygen (see Fig.5.10b). To compare, FeO/Pt(111) was also tested at
CO/O2=1:10 for the reference. Again, as in the case of CO/O2=1:5 (see Fig.5.6), FeO/Pt(111) has
shown the reaction rate higher than MnOx/Pt(111), namely more than twice.
Figure 5.9 Kinetics of CO oxidation over 1 ML MnOx film on Pt(111) at CO:O2=1:10. Reaction conditions: 5.45 mbar
CO+54.55 mbar O2, balanced by He to 1 bar; 450 K, heating rate 1 Ks-1 (to reach 450 K). A kinetic curve for 1 ML
FeO film on Pt(111) is given as a reference.
a
b
93
Figure 5.10 CO TPD (a) of 1 ML MnOx film on Pt(111) and its decomposition (b) after reaction with
CO/O2=1:10, β=3 Ks-1.
5.3.2. Oxygen states differentiation
Additional structural characterization of the similarly prepared oxide surfaces (including
multilayer films) was performed MAX-lab beamline I511 (Lund University, Sweden) [194].
Briefly, the preparation UHV chamber of the endstation was equipped with LEED, QMS, sputter
gun and a load lock for sample transfer without breaking vacuum in the main chamber. The
latter houses a hemispherical analyzer for high pressure x-ray photoelectron spectroscopy (HP-
XPS) and high pressure (HP) cell with a separate load lock that allows catalytic reactivity
measurements at near-ambient pressure conditions (up to 2 mbar). The synchrotron light after
the monochromator enters the cell via a Si3N4 membrane window while the electrons escape to
the analyzer through a 1mm aperture situated some millimeters in front of the sample. The
crystal is heated in the preparation chamber by direct electron bombardment of the sample
plate, while in the HP cell the heating is provided through electron bombardment of the high
pressure cell wall. The sample temperature is measured both by a type K thermocouple and a
pyrometer. The reaction was carried in a flow regime and monitored by QMS. Experimental
guidance and operation was performed and assisted by J. Weissenrieder and M. Soldemo,
members of Material Physics Division, KTH Royal Institute of Technology.
a
b
94
The evolution of O-species, present in O2-treated samples, was monitored in situ and ex
situ by XPS. Already at room temperature a shoulder on the high binding energy (BE) side of the
lattice oxygen peak at 529.2 eV was detected. The intensity of the shoulder increased
significantly at 450 K as shown in Fig.5.11. This indication of a newly formed O-species, active in
the CO oxidation reaction, is fully consistent with the results described above. The ‘as prepared’
film was of Mn3O4 stoichiometry as concluded from XPS O 1s normalization by O(2×2)-Pt(111)
and LEED pattern. After O2 treatment in the HP cell the original LEED pattern was gone, which
again certifying the amorphous state of the final O-enriched film. Oxygen content in the oxide
film increased roughly 1.3 times after O2-treatment in the HP cell. The binding energy of the
main line in the O 1s spectrum goes from approximately 529.2 eV to 529.4 eV after oxidation.
Figure 5.11 XPS spectra for O 1s (h =750 eV) of a monolayer MnOx film grown on Pt(111): as prepared,
after 10-5 mbar O2 in the HP cell at room temperature, after 1.3 mbar O2 in the HP cell at 450 K. Film was prepared
by 1 ML Mn deposition in 5·10-8 mbar O2 at room temperature followed by postoxidation in 10-6 mbar O2 at 850 K.
The higher binding energy component is shifted approximately 1.4 eV. At the same time (not
shown here) the Mn 2p strong line shifts to higher BE after oxidation from around 640.7 eV to
641.1 eV. A low binding energy shoulder at ~ 638.9 eV is almost, all indicative of a higher
95
oxidation state. The line profile of Mn 2p is rather complicated (as in several 3d elements) with
multiplet splitting complicating detailed analysis.
Fig.5.12 shows the kinetics of CO oxidation at different temperatures over 1 ML MnOx
film on Pt(111) upon reaching a steady state regime (at 450 K) at a constant reactants ratio
CO/O2=1:10. After 500 K the film experienced two effects: it undergoes deactivation caused by
carbon deposits what could be concluded from CO2 concentration decrease in the end of each
temperature step; the following rapid gain of the reactivity at 530 K could be caused by Pt(111)
surface opening - a typical mass transfer limit is comparable from QMS trends [194].
Figure 5.12 CO oxidation kinetics at different temperatures in flow mode over 1 ML MnOx/Pt(111) upon reaching a
steady state regime (450 K), CO/O2=1:10, p(HP cell) ~1 mbar. Trends are shifted for clarity (~50% CO is converted).
Fig.5.13 displays the reaction kinetics at varying CO partial pressure at 450 K while
keeping O2 partial pressure constant. It can clearly be seen that CO2 production proportionally
follows the change of CO concentration up to CO/O2 reaching 1:12; this suggests that the
reaction is first order for CO. In parallel the amount of CO2 in the gas phase over the oxide film
surface at 450 K and CO/O2=1:10 was sufficient to be detected by in situ HP-XPS (not shown).
Such CO pressure dependence can be definitely attributed to the catalytic reaction over the
surface of the MnOx film. It is worth to notice that in the similar experiment with stepwise
96
increasing CO partial pressure at fixed O2 pressure over clean Pt(111) no response like CO2
production following CO concentration change was observed. A 2 ML thick MnOx film appeared
to be less active in CO2 production. In the flow mode QMS signal 44 for 1 ML MnOx/Pt(111), 2
ML MnOx/Pt(111) and Pt(111) had the ratio 1:0.6:0.13 respectively.
Figure 5.13 CO oxidation kinetics at different CO partial pressures in flow mode over 1 ML MnOx film on
Pt(111) at 450 K and constant O2 partial pressure (O2 flowrate through the HP cell is 5 mL/min).
However, when increasing the CO concentration in O2 above 1:10 the monolayer film
cannot convert CO anymore and ‘dies’. This happen together with the disappearance of the
high BE side shoulder in the O 1s spectrum under CO-rich reaction conditions (the shoulder
cannot be restored by pure O2 treatment and the film remains inactive in a new reaction runs,
not shown here). This observation is in agreement with earlier experiments where the ratio of
the reactants was found to be crucial requirement for the film stable performance (1:10 against
1:5, Figs. 5.9 and 5.6 respectively). CO adsorption could be detected on the dead surface (in situ
HP-XPS, C 1s at around 286-287 eV). Therefore, the 'dead' MnOx surface is considered to be a
result of CO poisoning: manganese oxide gets reduced at the high CO partial pressure.
Moreover, the spent surface has a different line profile in Mn 2p than the active (a clear
doublet in the 2p3/2). The peaks are located at around 640.8 eV and 641.6 eV respectively (a
97
split of approximately 0.8 eV), what could certify a different structure or coordination to C. Such
a feature is reproducible for a 2 ML MnOx film as well.
CO can be removed from the 'dead' 1 ML surface by oxygen treatment, but the catalyst
will not regain its all native properties. There is some permanent change. In contrast the 2 ML
system cannot be made “active” again even though the CO can be removed during
regeneration.
5.4 Summary
Ultrathin MnOx films epitaxially grown on Pt(111) were studied in the low
temperature CO oxidation reaction at near-atmospheric pressures.
Under O-rich conditions initially crystalline monolayer MnOx films were
transformed into an amorphous MnO2 structure, assigned to O-Mn-O trilayer.
These films expose more weakly bound O-species as compared to those in
MnO(19×1)-Pt(111) films.
However, for stable catalytic performance MnOx/Pt(111) requires more oxidative
reaction conditions (CO/O2=1:10), otherwise the film dewets.
98
Chapter 6
General trends in CO oxidation over thin oxide films and
concluding remarks
6.1 General trends in CO oxidation over thin oxide films
Now, after all the films were introduced in previous sections, it is appropriate to
compare on their reactivity in attempts to understand what rational background stays behind
the reactivity of ultrathin oxide films on metal supports in low temperature CO oxidation at
elevated pressures. For this, we first present thermal desorption spectra of all studied films
after oxygen pretreatment on Fig.6.1. RuOx/Ru(0001), FeO/Pt(111) and MnOx/Pt(111) placed
together in the bottom portion of Fig.6.1 exhibit each well-pronounced evidence of weakly
bound oxygen species different from lattice oxygen in pristine films as indicated by dotted lines
for clarity: 420, 840 and 950 K correspondingly. RuOx/Pt(111) demonstrates a number of
broadened features around the 'main' oxygen peak what could be interpreted as a mixture of
phases other from RuO2(110). RuOx thin films prepared on Pt(111) did not show rectangular
(110) structure and demonstrated poor performance under reaction conditions in comparison
with that of RuO2(110)/Ru(0001). For comparison a spectrum of 2 MLE thick ZnO/Pt(111) is
given: this film is fully inactive under the same conditions.
Quantitatively the difference in catalytic performance can be described in terms of
desorption energies of weakly bound oxygen (WBO)-species. Using for the rough approximation
the Redhead formula (Eq.2.23), desorption energies were evaluated through the peak
temperatures 420, 825, 840, 950 and 1040 K (ZnO has only lattice oxygen) as 109, 219, 223, 253
and 283 kJ·mol-1 for RuOx/Ru(0001), RuOx/Pt(111), FeO/Pt(111), MnOx/Pt(111) and ZnO/Pt(111)
respectively. Qualitatively the values of WBO desorption energy are in line with activation
energies of CO oxidation: 58±4 and 113±5 kJ·mol-1 for RuOx/Ru(0001) [100] and FeO/Pt(111)
[14] correspondingly, measured under identical reaction conditions. When the reactivity of the
films is plotted against the corresponding desorption energies of WBO-species, the linear
99
dependence can be observed (see Fig.6.2). Therefore, the lower the desorption temperature of
WBO, the more active is the film in the reaction and vice versa. The ratio of WBO to lattice
oxygen is not that crucial as well as the origin of the film (deposited or native).
Figure 6.1 TPD spectra of molecular oxygen from decomposition in UHV of the films pretreated with O2 in the high
pressure cell (20 mbar O2, 450 K, 10 min). Heating rate β=3 Ks-1
Figure 6.2 The rate of the CO oxidation reaction as a function of desorption energy of weakly bound oxygen for
ultrathin films (grown on 2-side polished crystals, besides RuOx/Pt(111)). The rate is given only for the first 5
minutes of the reaction (zero conversion conditions). Desorption energy is calculated by the Redhead formula.
RuOx and ZnO are 4.5 MLE and 2 ML thick (closed at this coverage), while FeOx and MnOx are monolayer films.
100
The case of ZnO/Pt(111) is different: it is a catalyst with a particular reaction
mechanism. The film gets active in low temperature CO oxidation when it partially covers the
substrate surface, i.e. in submonolayer regime: as it was shown in Ch.4, the reaction occurs on
oxide-metal interface.
It is noteworthy to say that similar attempts to link thermodynamic and kinetic
parameters were made regarding such a particular case of oxidative heterogeneous catalysis as
the Deacon process (HCl oxidation) over rutile oxide surfaces. In the theoretical work with DFT
calculations [195] the authors found volcano dependences of TOFs in oxidation of hydrogen
halides as a function of the dissociative chemisorption energy of O2 and the adsorption energy
of OH on a variety of rutile (110) surfaces. Later, experimentally it was shown [196] that the
same is correct for bulk oxides.
Therefore, on ultrathin films grown on metals two CO oxidation scenarios (reaction
channels) are possible as schematically shown on Fig.6.3. Qualitatively it can be stated that flat
closed films act as oxygen exchangers due to their capability to activate oxygen from the gas
phase and therefore provide weakly bound oxygen (WBO) species under reaction conditions
(surface pathway, Fig.6.3a). Quantitatively the films reactivity follows the desorption energy of
top layer oxygen (Figs. 6.1 and 6.2). Films partially covering the metal substrate surface may
accommodate the reactants on the oxide-metal interface and there facilitate their interaction
(interface pathway, Fig.6.3b).
Figure 6.3 Scenarios of the CO oxidation reaction over ultrathin oxide films: a - surface pathway, CO2 is formed on
the film surface; b - interface pathway, the reaction happens on the metal-oxide interface.
a
b
101
6.2 Concluding remarks
Relying on the recent advances in the subject area, the present work comprises a broad
experimental basis to formulate basic concepts for low temperature CO oxidation over ultrathin
oxide films grown on metals. The performed research has established relationships between
the films structure and their reactivity under technologically relevant conditions with a
particular focus on thickness influence.
The key findings of the thesis can be summarized as follows.
A characteristic attribute of the closed films (RuOx/Ru(0001), MnOx/Pt(111) and studied
earlier FeO/Pt(111) by Y.-N. Sun [14]) as a one of the key observations in this thesis refers to
their ability to undergo oxygen enrichment at high chemical potential of oxygen in the CO
oxidation reaction under net oxidizing conditions. In that case the catalytic 'power' of a
certain film is defined by the strength of chemical bond between the transition metal of the
film and weakly bound oxygen species developed under reaction conditions.
Submonolayer films only partially covering the substrate surface open up another reaction
pathway at the oxide/metal boundary, which promotes the reactants activation and
interaction. The catalyst activity is defined by availability of the interface border, i.e. its
perimeter.
Oxygen binding energy can serve as a good descriptor for the CO oxidation reaction on
oxide films.
Thickness/coverage dependence allows one to determine the predominant reaction route.
When dealing with a new class of materials scientists highlight its typical properties and
categorize into subgroups with similar features. Scientific novelty of the present thesis is
addressed as a systematic approach when studying the reactivity of ultrathin oxide films.
Previously, the knowledge about thin films as model catalysts was quite segmental and diffuse
to give an overall understanding of their catalytic activity. Having accumulated my own results
in a separate cycle as a continuation of earlier works of our group, here I in a cooperation with
102
my colleagues suggest a universal way to 'handle' ultrathin oxide films in catalytic research and
to analyze and interpret data.
The practical importance of the performed work could be fully comprehended when
transferring its achievements to other more complicated chemical reactions over new kinds of
ultrathin oxide films as a good ground in this field. 'Playing' with CO oxidation over mono-
component films as a probe reaction should be logically replaced by catalytic processes of
higher demand, over multi-component ultrathin oxide films to bridge the 'timeless' gap
between industrial research and model studies in heterogeneous catalysis.
My personal contribution includes experimental work with all techniques overviewed in
Ch.2 (LEED, AES, TPD and GC) as well as data analysis and interpretation (Chs.3-6). The key
results and observations of the thesis were well-documented in several publications and
presented to the corresponding scientific community on dedicated conferences as it is shown in
App. A.
103
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List of abbreviations
AES Auger electron spectroscopy
AFM atomic force microscopy
DFT density functional theory
GC gas chromatography
GID grazing incidence diffraction
HP (cell) high pressure cell
HREELS high resolution electron energy loss
spectroscopy
HRTEM high resolution transmission
electron microscopy
HPXPS high pressure x-ray photoelectron
spectroscopy
IMFP inelastic mean free path
LEED low energy electron diffraction
LEIS low energy ion scattering
MEIS medium energy ion scattering
ML/MLE monolayer/monolayer equivalent
NEXAFS near edge x-ray absorption fine
structure
OCM oxidative coupling of methane
OMS octahedral molecular sieves
PLD pulsed layer deposition
PVD physical vapor deposition
QMB quartz microbalance
RBS Rutherford backscattering
SMSI strong metal-support interaction
SPM scanning probe microscopy
STM scanning tunneling microscopy
TCD thermal conductivity detector
TDS thermal desorption spectroscopy
TEM transmission electron microscopy
TMO transition metal oxide
TOF turnover frequency
TPD temperature-programmed desorption
TPR temperature-programmed reaction
UHV ultrahigh vacuum
VOCs volatile organic compounds
WGSR water gas shift reaction
XAS x-ray absorption spectroscopy
XRD x-ray diffraction
XPS x-ray photoelectron spectroscopy
118
List of figures and tables
Figure 1.1 Ultrathin FeO film grown on Pt(111)……………………………………………………………..3
Figure 2.1 Horizontal cross-sectional view of the UHV chamber……………………………………6
Figure 2.2 Sample holder and high-pressure cell sketch………………………………………………..7
Figure 2.3 Gas chromatograph setup…………………………………………………………………………….8
Figure 2.4 Diagram of a standard 4-grid LEED optics apparatus…………………………………….9
Figure 2.5 Universal curve of electron mean free path in solid matter [25]………………..11
Figure 2.6 Diffraction process………………………………………………………………………………………12
Figure 2.7 Ewald sphere synthesis for 2D case…………………………………………………………….13
Figure 2.8 Terminology for surface lattices………………………………………………………………….14
Figure 2.9 Crystallographic notation systems………………………………………………………………15
Figure 2.10 Auger electron emission mechanism and shells nomenclature…………………16
Figure 2.11 Energy spectrum of electrons coming from a surface irradiated with a beam of
primary electrons………………………………………………………………………………………………………..17
Figure 2.12 AES spectrum of ruthenium oxide as an example………………………………………17
Figure 2.13 AES - electron energy analysis……………………………………………………………………18
Figure 2.14 Optimal energy of a primary beam and yield of Auger electrons………………18
Figure 2.15 Schematic representation of QMS setup for TPD experiment……………………20
Figure 2.16 Potential energy diagrams for a diatomic molecule approaching and interacting with
a solid surface………………………………………………………………………………………………………………23
119
Figure 2.17 Schematic TPD spectrum……………………………………………………………………………23
Figure 2.18 Orders of desorption process. Sets of TPD profiles illustrate various desorption
kinetics………………………………………………………………………………………………………………….…….26
Figure 2.19 Complete analysis of TPD data [28]…………………………………………………………..28
Figure 2.20 Desorption process in terms of the transition state theory……………………….29
Figure 2.21 Operational principle of a general purpose gas chromatograph……………….31
Figure 2.22 Schematic drawing of a component peak on a chromatogram…………………32
Figure 2.23 Van Deemter plot………………………………………………………………………………………34
Figure 2.24 Representative flow through chromatographic columns………………………….35
Figure 2.25 The working principle of a thermal conductivity detector…………………………37
Figure 3.1 Real space geometry of Ru(0001) and O-adlayers unit cells [70]………………..40
Figure 3.2 Surface geometry of RuO2(110)……………………………………………………………………41
Figure 3.3 LEED patterns (70 eV) of different O-phases on Ru(0001)
used in this work………………………………………………………………………………………………………….42
Figure 3.4 Reaction kinetics for CO oxidation on clean Ru(0001) crystal, O-adlayers and thin
oxide films grown on Ru(0001)…………………………………………………………………………………….44
Figure 3.5 Sequential CO oxidation runs over the same
5 MLE thick RuOx film on Ru(0001)……………………………………………………………………………..45
Figure 3.6 Structure-reactivity dependence of CO2 production over RuOx films on Ru(0001),
O/Ru(0001) layers and initially clean Ru(0001) as a function of oxygen coverage measured
before the reaction……………………………….……………………………………………………………………..46
Figure 3.7 STM images of RuOx films grown on Ru(0001)………………………………………………47
120
Figure 3.8 TPD spectra of oxygen from decomposition in UHV of oxygen pretreated oxide films
RuOx/Ru(0001)…………………………………………………………………………………………………………….48
Figure 3.9 CO oxidation at different temperatures
over 5 MLE RuOx films on Ru(0001)…………………………………………………………………………...49
Figure 3.10 Pressure dependence of the CO oxidation reaction on 5 MLE RuOx films on
Ru(0001)………………………………………………………………………………………………………………………51
Figure 3.11 LEED patterns (70 eV) of an ordered 5 MLE RuO2(110)/Ru(0001) and a disordered
RuOx film on Ru(0001)……………………………………………………………………………………….……..…53
Figure 3.12 Effects of surface ordering on the reactivity
of thin RuOx films on Ru(0001)…………………………………………………………………………………...53
Figure 3.13 Epitaxial relationships of rutile RuO2(110) phase on a hexagonal surface Pt(111)
[106]……………………………………………………………………………………………………………………….…..54
Figure 3.14 LEED patterns (70 eV) of RuOx film on Pt(111) and clean Pt(111)………………55
Figure 3.15 CO TPD profiles from clean Pt(111) and RuOx film on Pt(111)…………………...55
Figure 3.16 Auger spectra of RuOx films of comparable thickness
on Ru(0001) and Pt(111)……………………………………………………………………………………………..56
Figure 3.17 Kinetics of CO oxidation over RuOx films grown on a one-side Pt(111) crystal and a
double-side Ru(0001) crystal………………………………………………………………………………………..57
Figure 3.18 TPD spectra of oxygen from decomposition in UHV of oxygen pretreated
RuOx/Pt(111) compared with RuOx/Ru(0001)…………………………………………………………….…57
Figure 3.19 Consistency between increase of the film reactivity and the absolute amount of
weakly bound oxygen (WBO) as a function of film thickness………………………………….……58
Figure 4.1 Structures of ZnO bulk material and thin film………………………………………………63
121
Figure 4.2 LEED patterns (at 60 eV) of ~1 MLE ZnO film on Pt(111), annealed in O2 at different
temperatures………………………………………………………………………………………………………………..65
Figure 4.3 CO TPD profiles from clean Pt(111) and ZnO films on Pt(111) annealed at different
temperatures in O2………………………………………………………………………………………………………..65
Figure 4.4 LEED patterns (at 60 eV) of ZnO films grown on Pt(111)
at the different film thickness…………………………………………………………………………………….….66
Figure 4.5 Auger spectra of ZnO films grown on Pt(111)
at the different film thickness……………………………………………………………………………………….66
Figure 4.6 ZnO/Pt(111) coverage calibration…………………………………………………….……………67
Figure 4.7 Auger signals of Pt and Zn as a function of the film thickness………………………..67
Figure 4.8 STM images of the “as prepared” ZnO films on Pt(111)……………………………….…68
Figure 4.9 CO oxidation over ZnO/Pt(111)………………………………………………………………….…..70
Figure 4.10 LEED patterns (60 eV) of the ZnO films
on Pt(111) before and after the reaction……………………………………………………………………...71
Figure 4.11 CO TPD for as prepared and post reacted surface of a ZnO film………………….71
Figure 4.12 LEED patterns (60 eV) of an alloyed Zn-Pt(111) surface before and after annealing to
600 K……………………………………………………………………………………………………………………….……..72
Figure 4.13 Auger spectra of an alloyed Zn-Pt(111) surface before and after reaction.….72
Figure 4.14 CO TPD for as prepared and post reacted surface of an alloy Zn-Pt(111)….….73
Figure 4.15 CO2 production measured during TPR over 1 MLE ZnO film on Pt(111)…..……73
Figure 4.16 LEED patterns (60 eV) of 1 MLE ZnO film on Pt(111) before and after TPR.….74
Figure 4.17 STM images of the “reacted” ZnO films…………………………………………………………75
122
Figure 4.18 CO TPD spectra of the 0.7 ML film before and after the CO oxidation reaction: film
dewetting……………………………………………………………………………………………………………………….76
Figure 4.19 Schematic representation of possible pathways of the reaction mechanism of CO
oxidation over ZnO/Pt(111) bilayer islands………………………………………………………………………78
Figure 4.20 TPD spectra of molecular oxygen from decomposition in UHV of ‘as prepared’ ZnO
films grown on two different Pt(111) crystals……………………………………………………………….…79
Figure 5.1 Schematic representation of bulk manganese oxides [145]…………………………….84
Table 5.1 Crystal structure of bulk manganese oxides……………………………………………………..85
Figure 5.2 Structural model of an oxidized MnO monolayer film on Pt(111) [193]………….87
Figure 5.3 LEED patterns of 1 ML MnOx films on Pt(111) at 70 eV……………………………………88
Figure 5.4 AES of 1 ML MnOx/Pt(111) subsequently oxidized through MnO, Mn3O4 and MnO2
steps…………………………………………………………………………………………………………………………………89
Figure 5.5 TPD profile of MnOx/Pt(111) pretreated with O2………………………………………….…90
Figure 5.6 Kinetics of CO oxidation over 1 ML MnOx film on Pt(111) at CO:O2=1:5……….…91
Figure 5.7 CO TPD of 1 ML MnOx film on Pt(111) and its decomposition after reaction with
CO/O2=1:5……………………………………………………………………………………………………………………….91
Figure 5.8 LEED patterns (70 eV) of a spent 1 ML MnOx film on Pt(111) after reaction CO/O2=1:5
and the subsequent first CO TPD run………………………………………………………………………………..92
Figure 5.9 Kinetics of CO oxidation over 1 ML MnOx film on Pt(111) at CO:O2=1:10…………92
Figure 5.10 CO TPD of 1 ML MnOx film on Pt(111) and its decomposition after reaction with
CO/O2=1:10……………………………………………………………………………………………………………………….93
123
Figure 5.11 XPS spectra for O 1s (750 eV) of a monolayer MnOx film grown on Pt(111): as
prepared and after O2 treatment……………………………………………………………………………………….94
Figure 5.12 CO oxidation kinetics at different temperatures in flow mode over 1 ML
MnOx/Pt(111)………………………………………………………………………………………………………………….…95
Figure 5.13 CO oxidation kinetics at different CO partial pressures in flow mode over 1 ML MnOx
film on Pt(111) at 450 K and constant O2 partial pressure………………………………………………...96
Figure 6.1 TPD spectra of molecular oxygen from decomposition in UHV of the films pretreated
with O2 in the high pressure cell……………………………………………………………………………………….99
Figure 6.2 The rate of the CO oxidation reaction as a function of desorption energy of weakly
bound oxygen for ultrathin films………………………………………………………………………………………99
Figure 6.3 Scenarios of the CO oxidation reaction over ultrathin oxide films……………………100
124
Appendix A: List of publications and
conferences attended
PUBLICATIONS
1. Y. Lei, M. Lewandowski, Y.-N. Sun, Y. Fujimori, Y. Martynova, I. M. N. Groot, R. Meyer, L.
Giordano, G. Pacchioni, J. Goniakowski, C. Noguera, S. Shaikhutdinov, and H.-J. Freund, CO + NO
vs CO + O2 reaction on monolayer FeO(111) films on Pt(111), ChemCatChem 3 (2011), p. 671-
674.
2. Y. Martynova, B. Yang, X. Yu, J. A. Boscoboinik, S. Shaikhutdinov and H.-J. Freund, Low
Temperature CO Oxidation on Ruthenium Oxide Thin Films at Near-Atmospheric Pressures,
Catal.Lett. 142 (2012), p.657-663
3. B. Yang, W. E. Kaden, X. Yu, J. A. Boscoboinik, Y. Martynova, L. Lichtenstein, M. Heyde, M.
Sterrer, R. Włodarczyk, M. Sierka, J. Sauer, S. Shaikhutdinov, H.-J. Freund, Thin silica films on
Ru(0001): Monolayer, bilayer and three-dimensional networks of [SiO4] tetrahedral, Phys. Chem.
Chem. Phys. 14 (2012), p.11344-11351
4. Y. Martynova, B.H. Liu, M.E. McBriarty, I.M.N. Groot, M. Bedzyk, S. Shaikhutdinov, H.-J.
Freund, CO oxidation over ZnO films on Pt(111) at near-atmospheric pressures, J. Catal. 301
(2013), p. 227-232.
5. Y. Martynova, M. Soldemo, J. Weissenrieder, S. Shaikhutdinov and H.-J. Freund, CO oxidation
on Mn-oxide thin films, (2013), in preparation
6. Y. Martynova, S. Shaikhutdinov, H.-J. Freund, CO oxidation on metal supported ultrathin
oxide films: what makes them active? (2013). ChemCatChem, submitted.
CONFERENCE APPEARANCES
1. European Conference on Surface Science (ECOSS-28), 28.8-2.9.2011, Wroclaw, Poland
Y. Martynova, S. Shaikhutdinov, H.-J. Freund
Ru(0001) vs RuOx thin films in low temperature CO oxidation (oral presentation)
125
2. European Conference on Surface Science (ECOSS-29), 3.9-7.9.2012, Edinburgh, Scotland, UK
Y. Martynova, B.H. Liu, M.E. McBriarty, I.M.N. Groot, S. Shaikhutdinov, H.-J. Freund
CO oxidation over ultrathin ZnO films on Pt(111) (poster)
3. 46. Jahrestreffen Deutscher Katalytiker, 13-15.3.2013, Weimar, Germany
Y. Martynova, S. Shaikhutdinov, H.-J. Freund
CO oxidation over metal supported ultrathin oxide films (poster)
126
Appendix B: Curriculum Vitae
Yulia Martynova
geboren am 23.10.1986
in Krasnokamensk, Russland
Schulausbildung
1993 -1998 Schule-Gymnasium № 9, Krasnokamensk, Russland
1998 - 2001 Lyzeum № 7, Tomsk, Russland
2001 - 2003 Lyzeum an der Polytechnischen Universität Tomsk, Russland
Studium
2003-2007 Polytechnische Universität Tomsk, Russland
2007 als Bachelor der Technik und Technologie abgeschlossen
2007-2008 Polytechnische Universität Tomsk
2008 als Diplom-Ingenieur in Chemischen Technologie
des Öls und des Gases abgeschlossen
Berufliche Tätigkeit
07/2008 – 10/2009 Laborantin bei NIOST, Labor der Katalysatoren und
katalytischer Systeme, Tomsk, Russland
Forschungsthema: Preparation and testing of new titanosilicate
catalysts for direct propylene epoxidation with oxygen
Promotion
02/2010 Doktorandin an der Technischen Universität Berlin und
Forschungstätigkeit unter der Leitung von
Herrn Prof. Hans-Joachim Freund am Fritz-Haber-Institut
der Max-Planck-Gesellschaft, Abteilung Chemische Physik,
Berlin, Deutschland
Titel der Dissertation: CO oxidation on metal supported
ultrathin oxide films
127
Erklärung an Eides statt
Ich versichere an Eides statt, dass ich die vorliegende Arbeit selbständig verfasst und nur
die angegebenen Hilfsmittel und Quellen benutzt habe.
Berlin, 28. März 2013
