Low-Dimensionality and Epitaxial Stabilization
in Metal Supported Oxide Nanostructures:
Mn O on Pd(100) Mn O
Cesare Franchini† and Francesco Allegretti‡
†
Faculty of Physics, Computational Materials Physics, Universität Wien and Center for
Computational Materials Science A-1090, Wien, Austria
‡
Institute of Physics, Surface and Interface Physics, Karl-Franzens University Graz, A-8010
Graz, Austria
We present a survey of the growth and structure of manganese oxide
nanolayers on a Pd(100) substrate, investigated in two different
thickness regimes through a plethora of surface science techniques
(STM, AFM, LEED, SPA-LEED, XPS, XAS, and HREELS) and
state-of-the-art theoretical tools (DFT and hybrid DFT approaches).
In the high thickness regime, above 10-15 monolayers, depending on
the preparation conditions different films with specific growth
direction and stoichiometry are formed. At low and intermediate
pressure (
1 10
mbar and 2-5 10
mbar, respectively) the
oxide structures are already bulk-like in terms of their in-plane
lattice constant and can be assigned to MnO(111) and MnO(100),
respectively. At high pressures (
5 10
mbar), Mn O
layers
(001)-oriented are obtained by oxidation of MnO(100). To explore
the epitaxial (geometric) relationships that favor the growth of the
2
different oxide phases, we have investigated the atomistic details of
different oxide/oxide and oxide/metal interfaces. In particular, we
have
addressed
the
issue
of
the
stability
of
the
Mn O (001)/MnO(001) interface and determined the phase
stability diagram of Mn O /Pd(100) phases at a Mn coverage of
about 1 ML. In the latter low thickness regime we have identified
nine different two-dimensional (2D) phases, which are novel in
terms of their structural and electronic properties. These nanophases
can be classified according to similar building block units and
described either as O-Mn-O MnO(111)-like trilayers or in terms of
metal-deficient MnO(100)-like monolayers, and therefore we argue
that they mediate the epitaxial growth of MnO thicker films on
Pd(100) by providing structurally graded interfaces. Moreover, the
formation of O or Mn vacancies drives the transition between 2D
phases with similar structural units but different lattice periodicity,
indicating
that
ion
vacancies,
mixed
valence
states
and
substoichiometry lie at the basis of the architectural flexibility in the
monolayer regime. Interestingly, the latter concepts play a major
role in the more complex class of functional oxides such as the
manganites, of which binary manganese oxides are the simplest
parent compounds.
3
10.1 Introduction
Since several decades, the study of transition metal oxides
represents a field of active and intense research. Thanks to the
uniquely rich spectrum of structural, electronic and magnetic
properties, which cover the range from metallic behavior to
magnetic insulators and encompass spectacular phenomena as
superconductivity, charge ordering and colossal magnetoresistance,
these compounds have attracted the interest of physicists, solid state
chemists and materials scientists. Renewed interest has been fuelled
by the advances in the synthesis processes of oxide materials, which
now allow control over the structure and stoichiometry at the level
of a unit cell. This has opened up new far-reaching perspectives for
both fundamental studies and technological applications, in
particular in the field of functional and multi-functional oxides,
many of which incorporate in their structure one or more transition
metal atom species. Functional oxide materials have unique physical
and chemical properties, which can be suitably controlled and
modified by means of external stimuli, such as changes of
temperature and pressure, electric or magnetic fields and the
adsorption of foreign atoms or molecules, thereby enabling the
development of new functionalities. As a result, a wide range of
intriguing
phenomena
can
be
observed,
ranging
from
superconductivity to ferroelectricity, piezoelectricity and exotic
magnetic behaviors. The ability to tailor and tune the properties of
these functional materials thus guarantees a high potential for
4
applications
in
micro-
and
nano-electronics,
spintronics,
heterogeneous catalysis, gas sensing, energy harvesting, etc.
In addition to these developments, novel properties can also
arise by the scaling down of the dimensions of oxide-based devices.
At the nanoscale, in fact, effects arising from the reduced
dimensionality, the geometric confinement, and the proximity of
surfaces and interfaces as well as effects arising from the coupling of
the oxide to the supporting material (so-called substrate-induced
effects, such as electronic hybridization, charge transfer and elastic
strain) may convey new physical and chemical properties and
therefore new functionalities to oxide systems. This potential is
indeed reflected in the enormous impulse given to fundamental and
applied research in the fields of nanoscience and nanotechnology.
Among low-dimensional nanostructured oxide systems, a prominent
role is played by transition metal oxide ultrathin films supported on
metal surfaces. Just to name a few technological applications, they
are typically employed as gate dielectrics and tunnelling barrier
layers in conventional and novel electronic devices [9, 63], as
protective layers in corrosion prevention and inhibition [42], as gas
sensor materials [34, 71], as support surfaces in the field of
heterogeneous catalysis [22]. In fundamental scientific studies, on
the other hand, thin films of transition metal oxides grown on metal
single crystal surfaces constitute preferred model systems for the
elucidation of emerging phenomena at the atomistic scale. Not only
the presence of the metal substrate allows circumventing charging
problems arising from the insulating character of many oxides, but
5
also it may lead - through the active participation in the elastic and
electronic coupling - to the stabilization of novel hybrid systems,
whose structural, electronic and magnetic properties bear no
correspondence to those of bulk oxides [60].
With this general frame in mind, in this chapter we focus on
the study of ultrathin manganese oxide films epitaxially grown on
Pd(100), providing a brief review of a combined theoretical and
experimental investigation performed by our research groups at the
University of Vienna and at the Karl-Franzens University of Graz.
All calculations presented in the present chapter have been
performed using the Vienna ab initio Simulation Package (VASP)
[46, 47] in the framework of density functional theory (DFT) [45]
and hybrid DFT [40].
Due to the ability of manganese to assume different
oxidation states, ranging from +2 to +7 [39], Mn oxides in bulk
compounds exhibit a number of stoichiometries with a complex
phase diagram [28]. Their architecture is dictated by the ability of
the Mn/O complexes to assemble by corner-sharing, edge-sharing or
double corner-sharing, such that more than thirty Mn oxide mineral
phases occur in nature [64]. As a consequence, Mn oxides display a
richness of behaviors which render them attractive systems in
applications of heterogeneous catalysis [10] (for example as electrocatalysts in fuel cells [72, 77]) as well as in applications as electrode
materials in solid state batteries [8] and in the environmental waste
treatment [21, 75]. Moreover, Mn oxides are parent compounds of a
particular class of functional oxides, the manganites, the properties
6
of which are determined by the complex interplay among spin,
orbital and lattice degrees of freedom, which results in outstanding
phenomena [66] such as giant magnetoresistance [76], metalinsulator transitions and orbital orderings[41]. These spectacular
properties promise to find application in the fields of magnetic
recording and novel spintronic devices. As to the study of Mn oxides
in reduced dimensions, an increasing effort has been devoted in the
last decade, which has revealed the richness of unusual behaviors
and an even higher degree of complexity relative to bulk Mn oxides.
For
example,
nanoparticles,
the
ferromagnetic
anomalous
with
behavior
respect
of
to
small
the
MnO
observed
antiferromagnetic ordering in the bulk phase, has been first predicted
by DFT [59] and then confirmed by subsequent experiments [50].
Recently, ferromagnetism of MnO and Mn O nanowires has also
been observed [56], whereas unconventional exchange bias
interaction has been reported for oxide coated Mn nanoparticles with
a Mn O shell [73], the latter effect being attributed to an unusual
spin alignment sequence at the interface.
To understand the unconventional properties of lowdimensional Mn oxides, fundamental studies on model systems are
desirable, which require a detailed knowledge of the oxide structures
at the atomistic level in the effort to elucidate the structureproperties relationship. So far, these studies on model systems have
mainly focused on the growth of MnO thin films on the Ag(001)
surface. Due to the relatively large in-plane lattice constant of
MnO(001) (
=3.14 Å), the template has been chosen to ensure a
7
reasonable overlayer-substrate matching (
=2.89 Å), which is
expected to favor better epitaxial growth. Indeed, despite the still
large mismatch (about 9%) films of good quality have been obtained
[55, 74]. Further investigations cast light on the evolution of the
structural and electronic properties with the oxide thickness due to
the partial release of the epitaxial strain [57, 58, 19]. However, no
detailed studies of the oxide-Ag(001) interface in terms of the Mn
oxide phases formed at coverage below or about 1 monolayer (ML)
are available in the literature. Interface-stabilized Mn O
monolayers have instead been reported on two different substrates,
Rh(100) [61] and Pt(111) [38], but the literature is still sparse and
the unambiguous structural assignment of the oxide phases has not
been yet accomplished. In this context, we have recently performed
a joint experimental-theoretical investigation aimed to extensively
characterize the growth and structure of Mn O
ultrathin films on a
Pd(100) substrate. The film thickness range extended from a few Å,
corresponding to the interface-stabilized phases of the monolayer
regime, up to 30-50 Å, where bulk-like behavior sets in. The
detailed theory-experiment comparison for the Mn oxide systems,
which will be presented in the next section, proved a decisive factor
for the unambiguous assignment of the oxide phases, and it enabled
to identify the interface-stabilized oxide structures that mediate the
epitaxial growth. Moreover, it proved highly beneficial, in that it
allowed assessing the reliability of density functional theory (DFT)
and post-DFT approaches applied to the Mn oxides. This is a very
8
important issue in the light of the well-known difficulties
encountered by standard DFT in dealing with strongly correlated
materials [1].
10.2 Growth of Mn O
Layers on Pd(100)
The growth of Mn oxides on Pd(100) has been the subject
of a series of recent papers [3, 16, 4, 51, 26, 27], which explore in a
rather systematic way the phase stability diagram at different oxide
thicknesses. In particular, it has been shown that below 1 ML a
complex surface phase diagram with a multitude of novel Mn oxide
structures develops, the oxide structures being characterized by
specific structural building blocks and vibrational properties [51, 26,
27]. In the high thickness regime, upon deposition of 20-30 ML,
MnO(100) films exhibiting good long range order can be grown
epitaxially [3], despite the considerable lattice mismatch with the
substrate ( 14%). The MnO(100) structure can be preferentially
converted either into MnO(111) [3, 4] or Mn O (001) [16]
depending on the combination of temperature and applied oxygen
pressure. The evolution of the physical properties of the manganese
oxide layers on Pd(100) with respect to the changes in film thickness
provides a unique method to separate surface (2D) and bulk (3D)
effects. In fact, to a decrease of film thicknesses from the high
coverage (multilayer) to the low coverage ((sub)monolayer) regime
9
corresponds an increase of the surface to bulk ratio, and low
dimensional effects become more distinguishable. The fundamental
characteristics of the Mn O /Pd(100) system are discussed in
details in Sec. 2.1 and 2.2 and can be summarized as follows:
Mn O /Pd(100)
Low coverage regime,
0.75 monolayer (see Sec. 2.1). Ultrathin
layers of variable Mn O
stoichiometry, only 1-2 ML thick, can be
formed at the metal-MnO interface. At least nine different 2D
Mn O
phases on Pd(100) are found, which are novel in terms of
the known Mn oxide bulk crystal structures and which are stabilized
by the metal-oxide interface and by the confinement in the direction
perpendicular to the surface. These low-dimensional interfacestabilized phases may mediate the epitaxial growth of thicker layers
by providing structurally graded interfaces.
High coverage regime,
20-30 layers (see Sec. 2.2). MnO(100)
with bulk-like in-plane lattice constant is stable in a wide range of
pressure and temperature. Through the appropriate tuning of
temperature and oxygen pressure the MnO(100) films can be
transformed into MnO(111) (annealing at elevated temperatures or
reactive evaporation at lower oxygen pressures [
mbar]) or transformed into Mn O (001) surfaces (high temperature
oxidation at relatively high oxygen pressure [
10
mbar]).
10.2.1 Low Coverage Regime
The experimental phase stability diagram of Mn oxides on
Pd(100) below 1ML is depicted schematically in Fig. 1, where the
various Mn oxide nanolayer phases are ordered as a function of the
oxygen pressure
and of the oxygen chemical potential
during
the preparation procedure, and they are represented by their
corresponding scanning tunneling microscopy (STM) profile.
Fig. 1 Experimental schematic phase stability diagram of the interfacial Mn
oxides, presented as a function of the oxygen pressure
chemical potential
and of the oxygen
. The nominal coverage of Mn on Pd(100) is
0.75 ML.
From Ref. [51] with permission
The manifold Mn oxide phase diagram comprises nine
different nanophases which are characterized by specific windows in
the parameter space of the thermodynamic variables temperature (in
the total range: 600-800 K) and oxygen pressure (in the total range:
5 10
-5 10
mbar). We distinguish:
1. The oxygen-rich regime (5 10
mbar >
< 5 10
mbar)
(a) Two hexagonal phases (HEX-I and HEX-II), which are both
11
obtained at high
pressures
2. The oxygen-intermediate regime (5 10
mbar >
< 1 10
mbar)
(a) A c(4 2) structure and a stripe phase described as a uni-axially
compressed c(4 2), which are both stabilized at intermediate
pressures
(b) Two structures which were called chevrons (CHEV-I and
CHEV-II), because of their STM appearance
3. The oxygen-poor regime: (1 10
mbar >
< 5 10
mbar)
(a) Two reduced phases with complex structures, named waves and
labyrinth
(b) At the most reducing conditions, a third hexagonal phase
(HEX-III), commensurate with the Pd(100) substrate ( 2) along one
of the two
directions.
The extraordinary architectural flexibility of the interfacial
Mn oxides on Pd(100) and the electronic properties of the novel 2D
phases can be rationalized and understood through a synergic
combination of experimental [STM, low energy electron diffraction
12
(LEED),
high-resolution
electron
energy
loss
spectroscopy
(HREELS) and X-ray photoemission spectroscopy (XPS)] and
computational (DFT and hybrid functionals) techniques. Indeed, in
modern surface science theoretical calculations have become an
efficient complementing tool to the experimental observations.
Theoretical models based on educated guesses of possible structures
can be tested and directly compared with the experiments in order to
clarify the structural aspects and provide an atomistic interpretation
of the measured properties. In the present case, the joint
experimental-theoretical analysis reveals that the two oxygen-rich
phases (HEX-I and HEX-II) can be described in terms of O-Mn-O
trilayers with MnO(111)-like structures (see Sec.2.1.1), whereas the
intermediate oxygen regime (c(4 2) and chevrons) is based on a
compressed MnO(100)-like monolayer model (see Sec.2.1.2). At
low oxygen chemical potentials the waves and labyrinth structures
show in STM very complex unit cells whose link to the other phases
is less clear but seems still to be related to a MnO(100)-like wetting
layer, while the additional hexagonal phase HEX-III is of uncertain
attribution (see Sec.2.1.3).
Fig. 2 HREELS phonon spectra of the four Mn oxide submonolayer phases at high
and intermediate oxygen partial pressures. From the top to the bottom: HEX-I,
HEX-II, c(4 2) and CHEV-I. The statistical uncertainty in the peak position is
0.5 meV. For every Mn oxide phase, measurements performed on samples freshly
prepared in different days agree within 1 meV. Adopted from Ref. [51]
13
The recurrent MnO(100)- and MnO(111)-like structural
features are reflected in the phonon-loss spectrum displayed in Fig.
3(b)@. As expected, the common building blocks shared between
the different oxide phases result in similar phonon losses. All spectra
exhibit a clear single peak structure. The phonon loss is centered at
around 70 meV for the two HEX-I and HEX-II phases and shifts
down to 44-45 meV for the c(4 2) and CHEV-I phases.
Interestingly, the phonon spectra (not shown) of the CHEV-II,
waves and HEX-III phases, which are obtained by further lowering
the O chemical potential, are also characterized by a single peak at
43-45 meV. These findings provide clear indication of two distinct
regimes, a MnO(111)-like regime comprising phases with a higher
energy phonon loss (70 meV) and a MnO(100)-like regime with a
single phonon loss around 44 meV.
10.2.1.1 MnO(111)-like Phases (Oxygen-Rich Regime)
In this pressure regime two Mn oxide phases have been
detected, HEX-I and HEX-II, which are characterized by hexagonal
or quasi-hexagonal (i.e. distorted) symmetry linked to that of
MnO(111), and which display a phonon spectrum with a single loss
peak centered at around 70 meV.
HEX-I
For the sake of clarity, it is instructive to recall the structure of
MnO(111), which is made out of alternating O and Mn hexagonal
14
layers with ABC stacking sequence, each layer having in-plane
lattice constant
=3.14 Å [see Fig.(a)].
Fig. 3 (a) Real and (b) reciprocal lattice for a (1 1) MnO(111)-like hexagonal
structure on a Pd(100) surface. The lattice parameter of the overlayer is assumed
to be the measured bulk value
= 3.14 Å, whereas
=2.75 Å. In (a) only
one of the two symmetry domains is reported for clarity. (c) SPA-LEED twodimensional pattern measured at E=90 eV, for the undistorted HEX-I phase. (d)
LEED pattern of the distorted HEX-I phase, recorded at E=96 eV. (e) Real and (f)
reciprocal lattice associated with the distorted hexagonal MnO(111)-like phase.
Only one symmetry domain is shown in (e) for clarity, for which the lattice
parameter in the direction of the distortion ([011]
) is
=2.94 Å. In total,
four symmetry domains contribute to the reciprocal lattice in (f): two are obtained
from panel (e) with
at either +60 or -60 from
are obtained by rotating
by 90
, and the other two
relative to the substrate mesh. (g) Hard
sphere model simulating in the real space the moiré pattern that originates from
the interference of the quasi-hexagonal oxide lattice of panel (e) with the square
mesh of the Pd(100) substrate. Figure adapted from Ref. [26] and Ref. [51] with
permission
A perfect MnO(111) film would give rise to a hexagonal
reciprocal space pattern with lattice parameter
/3
=2.30 Å
.
The latter value is very close to the reciprocal space lattice
parameter of Pd(100) (
/
=2.28 Å
). The resulting LEED
pattern averaged over two hexagonal domains rotated by 90 , which
account for the different symmetry of substrate and overlayer, is a
circular array of 12 extra-spots superimposed on the (10) spots of the
15
substrate, as illustrated in Fig. (b). Although such a LEED pattern
has been indeed observed in the experiments, as demonstrated in
Fig.(c), the structural details sensibly depend on the preparation
conditions. It appears that this
hexagonal structure is
favored at slightly lower coverage (0.5-0.6 ML) or after repeated
oxidation cycles. In contrast, at 0.75 ML and after a single oxidation
step at high pressure the LEED pattern reported in Fig.(d) is
reproducibly observed: it still reflects a hexagonal-like lattice, but
the characteristic elongation of the overlayer spots indicates a clear
distortion of the ``perfect'' (1 1) structure into a
hexagonal structure. With respect to the perfect hexagonal lattice,
the
LEED pattern is distinctive of an incommensurate
overlayer with a contracted lattice vector
[011]
= 2.94 Å along
[Fig. (e)]. The simulated reciprocal space pattern of the
structure, shown in Fig. (f), reproduces very well the
experimental LEED pattern [Fig. (d)]. Interestingly, the comparison
between the
and
real space lattices of Figs.
(a) and (e) suggests that the better matching between overlayer and
substrate might be the driving force for the distortion: the initial
mismatch of 14% along the [011]
(
)/
direction is in fact reduced to
=7% upon distortion. This better epitaxial
relationship is likely to be at the origin of the increased stability and
reproducibility of the distorted HEX-I phase.
16
Fig. 4 (a-b) Top view of the geometrical models for
(a) and
(b) HEX-I MnO(111)-like trilayer phases: red spheres: O atoms;
light gray spheres: Mn atoms. Dashed lines delimit the 2D unit cell. (c-e)
Experimental (c,d) and simulated (e) STM images of the HEX-I phase;
Experimental images: sample bias U = +0.5 V (c), +0.6 V (d); tunneling current I
=0.13 nA (c), 0.15 nA (d). The simulated STM image has been calculated
considering tunneling into empty states between 0 and +0.5 eV. Taken from Ref.
[26] with permission
In passing, we note that the
model explains also
the moiré pattern observed experimentally in the STM images
[Fig.2(c)]. The latter results from the interference of the quasihexagonal Mn oxide overlayer with the square mesh of the Pd(100)
surface. Taking a ratio
/
16/15, a moiré pattern can be
generated with a geometrical model that displays modulations in
form of broad lines inclined with respect to the [011]
direction
and with an average periodicity of about 22 Å parallel to [011]
[see Fig. (g)]. In this model, only one type of overlayer atoms is
considered for simplicity, and a different color gradation (white,
grey, black) is used to emphasize the different lateral registry of
these overlayer atoms relative to the underlying matrix of substrate
atoms. For on-top/bridge location, the white/black color of the
overlayer atoms reflects the different height above the surface. The
remarkable similarity between the modeled moiré pattern [Fig. (g)]
and the experimental STM image [Fig.2(c)] suggests implicitly that
the HEX-I phase may consist of alternately stacked layers with
quasi-hexagonal symmetry containing only one single atom species,
17
as it is realized in the MnO(111) structure, which is therefore the
natural toy-model for the computational analysis.
The HEX-I model used in the DFT analysis is constructed by
cleaving the MnO structure perpendicular to the [111] direction to
form the smallest block containing 1 ML of Mn atoms in the O-MnO stacking, with a formal stoichiometry of MnO . A rigid sphere
sketch of the MnO
trilayer is given in Fig. 2(a), along with the
corresponding
the lattice constant
trilayer [Fig. 2(b)] obtained by shrinking
by 7%. Due to the free-standing (i.e.:
unsupported trilayers) computational setup, which is unavoidable
because of the incommensurate registry of the substrate/overlayer
system, the optimized 2D lattice constants are significantly smaller
than the measured one. Energetic considerations on the computed
relative stability of the two HEX-I models reveal that although the
perfect hexagonal structure is the most favorable, only a low energy
cost of 70 meV/cell is required to convert it into a
hexagonal structure. Therefore, the interaction with the substrate,
neglected in the unsupported trilayer model, might easily provide the
energy gain to reverse this energy sequence through the better
matching discussed above.
The soundness of the MnO trilayer model is confirmed by a
direct comparison of the HREELS and STM data with the computed
counterparts. Panels (d) and (e) of Fig. 2 show the comparison
between the measured and simulated atomically-resolved STM
images. Despite the low resolution of the experimental map the
18
agreement is acceptable, in particular in that the arrangement of
large bright protrusions is fairly reproduced; moreover, the
theoretical data allow assigning the bright spots to the topmost
oxygen atoms. Further convincing support for the proposed model
comes from the comparison between the experimental (70.5 meV)
and computed (
73 meV) distinct single dipole-active phonon
mode, which is associated to the vertical antiphase vibration of O
and Mn layers [26]. Finally, it is also worth noting that this trilayer
model is well compatible with the O
core level photoemission
data reported in Ref. [51], which show a strong component at 529.1
eV attributable to the surface O layer and a weaker line chemically
shifted to higher binding energy presumably related to the O layer at
the interface with the Pd(100) substrate. Indeed, a model consisting
of a MnO(111)-like bilayer, with only one layer of chemically
equivalent O atoms either at the interface or at the surface to yield a
formal MnO stoichiometry, is not only incompatible with the
photoemission data but it is also shown to be much less stable than
the corresponding trilayer model in a broad window of the O
chemical potential (Ref. [26]).
HEX-II
The O-Mn-O MnO(111) trilayer model is the crucial key for the
structural understanding of the HEX-II phase, which is obtained
from the HEX-I phase upon reduction of the oxygen partial pressure.
Unlike HEX-I, the STM appearance of HEX-II is characterized by a
19
well-ordered hexagonal array of triangularly shaped maxima [Fig.
3(c)], and the lattice constant derived from the STM and LEED
measurements [Fig. 3(a)] is
Å , corresponding to a
compressed (2 2) reconstruction of the MnO(111) termination. The
DFT
analysis
reveals
that
the
so-called
O-terminated
reconstruction displays the lowest surface energy in a wide range of
oxygen partial pressures [26]. The
structure is obtained by
removing one topmost oxygen atom per (2 2) unit cell from the
HEX-I structure, resulting in a
trilayer with in-plane
hexagonal lattice constant of 5.95 Å, as shown in Fig. 3(b).
Fig. 5 (a) LEED pattern of the HEX-II phase (E = 80 eV). (b) Top view of the
-
HEX-II model: red spheres: O atoms; light gray spheres: Mn atoms.
Dashed lines delimit the 2D unit cell, whereas triangles highlight the octopolar
Mn-O
pyramids and the O
trimers of the
simulated STM images of the hexagonal
termination. (c) Experimental and
HEX-II phase; sample bias U = +0.75
V; tunneling current I =0.1 nA. The simulated STM image has been calculated
considering tunneling into empty states between 0 and +0.75 eV. Unit cells are
indicated by full lines. Figure adapted from Ref. [26] and Ref. [51] with
permission
Compelling evidence for the O-terminated
model comes
from the comparison between the simulated STM and HREELS data
and the corresponding experimental results. The hexagonal pattern
of triangular bright protrusions in the experimental STM picture is
well reproduced by the simulated STM image [Fig. 3(c)], and the
20
triangular protrusions are therefore ascribed to the oxygen trimers
outlined in the sketch of Fig. 3(b). As for the HREELS spectrum, in
analogy with the precursor HEX-I phase, the HEX-II
phase
displays a single phonon loss peak centered at 69.5 meV [Fig.
3(b)@], characteristic of the MnO(111)-like structures. The
simulation reproduces well the position of the vibrational peak
(within 1-2 eV) and ascribes it to the concerted movement of the
surface oxygen trimers against the manganese atoms underneath.
10.2.1.2 MnO(100)-like Phases (Intermediate Oxygen Regime)
At intermediate oxygen pressures four different Mn O
phases
have been observed. At the high pressure end of this regime two
phases coexist (see Fig. 1): a quasi-hexagonal
phase
characterized by bands of parallel stripes separated by
(2
5.5 Å
), and the c(4 2) structure, commensurate with the Pd(100)
substrate as expressed by the transformation matrix M = [2,-1 / 2,1].
At the low pressure end, instead, the c(4 2) phase is often observed
in coexistence with the so-called chevron structures, chevron I and
chevron II, corresponding to the transformation matrices M = [2,-1 /
3,2] and [2,-1 / 5,3]. The phase boundaries between the two
chevrons structures and the c(4 2) structure are smooth and
continuous, suggesting that from the structural point of view these
phases are closely related. Indeed, the unit cells of the chevrons
structures can be generated from the c(4 2) structure by simply
21
extending the
unit cell vectors to adjacent antiphase positions as
illustrated in the geometric models of Fig. 2.1.2(b). In the following
the structural analogies of these phases are rationalized in terms of
MnO(100)-like monolayers.
Fig. 6 (a) LEED patterns, (b) real lattice models, and (c) reciprocal lattice
patterns of the c(4 2) (left panel), chevron I (middle panel) and chevron II
(right panel) structures. The respective LEED energies in (a) are 116, 108 and 60
eV. From Ref. [51]
c(4 2)-Mn O /Pd(100)
The LEED pattern of the c(4 2) structure, the central phase in this
pressure regime, is displayed in Fig. 2.1.2(a) and compared with the
simulated reciprocal space pattern of panel (c), which corresponds to
the real space periodicity given in panel (b). The structural model
that explains the c(4 2) Mn O
phase is depicted in Fig. 4. It is
constructed by placing a compressed MnO(100) monolayer on top of
Pd(100) and by creating a regular rhombic array of Mn vacancies, in
close analogy with the Ni/Co oxide monolayer phases, c(4 2)Ni O /Pd(100)
and
c(4 2)-Co O /Pd(100),
respectively,
obtained upon reactive evaporation of nickel/cobalt on Pd(100) [2,
5]. It is, however, to be noted that the lattice mismatch between
MnO (14%) and the substrate is much larger than that for NiO (7%)
and CoO (9.5%). The lateral compression required to fit a single
22
layer of the bulk MnO(100) structure onto the Pd(100) substrate and
the consequent formation of the Mn vacancies dominate the
structural rearrangement and determine the occurrence of two
distinct types of Mn atoms with different environment, as illustrated
in Fig.4(a): the Mn1 atoms are localized between two nearest
neighbor vacancies, while the Mn2 atoms form zig-zag Mn-Mn
chains along the [01 ] direction. Large relaxations take place in the
Mn O
layer, which is found to be well separated from the
substrate by z 2.3 Å, approximately 20% larger than the Pd bulk
interlayer distance. The lateral compression is accompanied by a
significant rumpling of the monolayer, yielding a corrugation of
0.23 Å, whereas the Mn vacancy breaks the local Mn-O bonds and
induces a planar outward relaxation
of the oxygen atoms, which
move closer to Mn1, resulting in the waving short-ranged ordering
highlighted in Fig.4(a).
The comparison with the experimental data, given in
Fig.4(b)-(c), validates the proposed c(4 2)-Mn O
structural
model. The measured STM image [inset of panel (b) in Fig.4] is
characterized by a rhombic arrangement of bright spots and dark
depressions connected by weaker segments. These features are very
well reproduced by the simulated STM picture and can be
interpreted in the following way: the black depressions reflect the
network of Mn vacancies, the single bright spots can be assigned to
the Mn1 species, whereas the light segments correspond to the zigzag Mn2-Mn2 chains embedded in the regular array of oxygen
23
atoms, which are not seen in the STM image. Further support for the
c(4 2)-Mn O
model is provided by the comparison of the
calculated phonon spectrum with the HREELS measurements
[Fig.4(c)]. The single peak structure centered at
43 meV revealed
by HREELS is very well reproduced by the hybrid functional
calculations (standard PBE, instead, underestimates the phonon
energy by
7 meV). The diagonalization of the dynamical matrix
enables to assign this specific phonon mode to the collective vertical
vibration of the oxygen sublattice against the Mn sublattice, a
vibrational mode typical of the ideal MnO(100) surface [26, 38].
Fig. 7 (a) Optimized structural model for the c(4 2)-Mn O /Pd(100) system.
The Pd(100) substrate is displayed with light gray large spheres, whereas Mn and
O atoms are depicted with small blue and pink spheres. Two distinct Mn species
are distinguishable: Mn1, sandwiched between two vacancies, and two Mn2 atoms
forming the zig-zag arrangements highlighted by the horizontal lines. Atoms
surrounding the Mn vacancies experience considerable strain (indicated by
arrows). (b) Comparison between simulated and experimental (inset) STM images
for the c(4 2)-Mn O /Pd(100) phase. Experimental parameters: sample bias U
= +0.8 V, tunneling current I= 0.25 nA.(c) Comparison between the measured
HREELS phonon value (vertical bar) and PBE (dashed line) and HSE (full line)
predicted dipole active modes. Adapted from Ref. [26] and Ref. [27]
Chevron structures
The chevron structures are adjacent to the c(4 2) phase in the phase
24
stability diagram and are obtained by lowering the O
partial
pressure. These phases, the name of which originates from the
``chevron-like'' STM motifs, can coexist locally with the c(4 2)
phase, as shown in Fig.5(a). Due to the close link with the c(4 2)
phase, presented in a pictorial fashion in Fig. 2.1.2 (b), the chevron
structures can be rationalized in terms of structural models inspired
by the c(4 2) structural features and by invoking the concept of
vacancy propagation. As discussed above, the periodicity of the
c(4 2) phase is given by the Mn vacancies distributed in a regular
array with rhombic unit cell. Therefore, within this picture and based
on Fig. 2.1.2(b), the chevron I and II structures may be naturally
obtained from the c(4 2) by propagating a Mn vacancy to a
neighboring lattice site. Accordingly, the stoichiometry becomes
and
, respectively, i.e. closer to that of MnO.
Compared to the c(4 2) phase, the chevron structures exhibit a
much higher degree of long range order, with large defect-free areas
extending over typical distances of 200 Å and interrupted by ordered
domain boundaries, which convey the characteristic ``chevron-like''
appearance [see Fig. 5(a)].
Fig. 8 (a) STM image showing the characteristic appearance of the chevron
structures and the local coexistence of the c(4 2), chevron I and chevron II
phases (140 Å
100 Å ; U = +1.0 V; I = 0.2 nA). (b) Top view of the geometrical
model for the Pd(100) supported chevron I phase: red spheres: O atoms; light
gray spheres: Mn atoms; green small spheres: Pd atoms underneath. Dashed lines
25
delimit the 2D unit cell and circles denote the position of the vacancies. (c)
Experimental and simulated STM images of the chevron I phase; sample bias U =
+1.0 V; tunneling current I =0.13 nA. The simulated STM image has been
calculated considering tunneling into empty states between 0 and +1.0 eV. Unit
cells are indicated by full lines. Figure adapted from Ref. [26]
The structural model for the chevron I structure, obtained by
removing 14% of Mn atoms from the compressed MnO(100)
monolayer, is depicted in Fig. 5(b). The geometrical optimization
yields again an oxygen terminated surface with a corrugation of 0.33
Å, and in analogy with the c(4 2)-
structure a moderate
displacement (0.13 Å) of Mn atoms towards the vacancy is found,
indicated by the arrows in Fig. 5(b). Figure 5(c) shows a comparison
between the experimental and simulated STM images. Although the
experimental resolution does not allow a detailed atomic comparison
between theory and experiment, the overall agreement is
satisfactory. The dark circular depressions in the images represent
the position of the Mn vacancies and determine the 2D unit cell of
this phase, as outlined by the full lines. The bright protrusions are
clearly due to manganese atoms, suggesting that electronic effects
contribute predominantly to the STM topography. Two distinct
bright features are observed in the experimental picture, one large
spot in the center of the unit cell and weaker flecks at the edges. The
simulated STM image together with the optimized geometrical data
allows the atomically resolved identification of these two kinds of
features. The manganese sublayer is itself slightly corrugated, due to
26
a small upward shift (0.05 Å) of the Mn atoms laying closer to the
vacancy and aligned in the [01 ] direction. These atoms, indicated
by filled dots in Fig. 5(b), are responsible for the two large bright
spots in the simulated STM image, which merge together in the
experimental image thereby giving rise to the wavy lines of bright
protrusions. The remaining spots (four per unit cell) arise from the
lower manganese atoms. In the experimental STM image their
intensity partially mixes up with the topmost Mn spots thus
contributing to the large bright feature centered in the cell and
partially providing the weaker spots along the edges. Finally, the
DFT-based analysis establishes that the
chevron I structure
is characterized by a single dipole-active phonon mode at about 40
meV, in good agreement with the HREELS measurements (phonon
loss at 44.5 meV). This peak reflects the vibrational fingerprint of
the parent MnO(100) not only in terms of the phonon loss energy,
very similar to the c(4 2)-
value (43.5 meV, see Fig. 3(b)@),
but also in terms of the atomic displacements producing this
vibration, which are associated to the vertical opposite movement of
manganese and oxygen atoms.
10.2.1.3 The Reduced Phases (Oxygen-Poor Regime)
There is a smooth transition from the c(4 2) and chevron structures
of the intermediate oxygen pressure regime to the structures in the
oxygen-poor regime denoted as labyrinth and waves. These two
phases are named according to their STM pattern shown in Fig. 1.
27
Although the complexity of these structures is a major obstacle for
the computational machinery, some fundamental properties can be
deduced from a careful analysis of STM and LEED data (not shown)
[51]. The labyrinth structure is described by the lattice vectors
13.5
0.5 Å and
between lattice vectors
21
= 11.0
and
0.5 Å,
2
( : angle
), the unit cell being rotated by
=
2 with respect to the Pd [011] direction. The unit cell of the
waves structure is characterized by
Å,
= 85
=
= 90
= 13.8
0.2 Å,
= 34
1
2 , which corresponds approximately to a (5 12)
superstructure with respect to the Pd(100) substrate. On the basis of
the measured phonon loss spectrum and of the occurrence of ordered
phase boundaries with the c(4 2) and chevron structures, these two
reduced phases with complex appearance are tentatively assigned to
MnO(100)-like wetting layers.
Finally, at the most reducing conditions a different structure
is observed, the HEX-III structure, shown on the left side of the
phase stability diagram of Fig. 1. The HEX-III structure is described
by the transformation matrix M = [0, 2 /
=[
(
, -1 /
, 1] (or equivalently M
, 1]) with respect to Pd(100), but corresponds also to a
)R30 superstructure with respect to a MnO(111) surface.
The last observation suggests that this phase might be related to a
MnO(111)-like model, possibly a bilayer with 1/3 of the surface
atoms removed. On the other hand, the matrix notation shows that
the periodicity of the HEX-III phase can be also obtained upon a
uni-directional 14%-compression of the c(4 2) structure. As a
28
result, a (7 2) coincidence lattice relative to the Pd(100) surface is
observed. This interpretation would be compatible with the very
similar phonon loss spectra of the HEX-III and c(4 2) phases, and
with certain features of the STM appearance at particular bias
voltages (indeed, the c(4 2) periodicity can itself be regarded as due
to a distorted hexagonal lattice of vacancies) [6]. This further
compression could hypothetically produce a more pronounced
rumpling of the corrugated monolayer, possibly resulting in a welldefined bilayer. The precise assignment of the HEX-III requires
detailed structural investigations, and to this purpose DFT
calculations are underway. It is, in any case, possible that the HEXIII structure is an interfacial phase that mediates the growth of
MnO(111)-oriented films with hexagonal symmetry, which will be
discussed in the next section.
10.2.2 High Coverage Regime
The reactive evaporation of 20-30 ML Mn on Pd(100) at moderate
temperature ( 620 K) and at an oxygen pressure in the range
between 2 10
and 5 10
(100)-oriented MnO films.
mbar leads to the stabilization of
The formation
of well-ordered
MnO(100) is evident from the spot profile analysis LEED (SPALEED) image presented in Fig. 2.2(a), which displays sharp reflexes
arranged in a square pattern and superimposed on a very low
background. The surface lattice constant determined from the
29
separation of the LEED spots measures 3.14 0.03 Å, identical to
the (100) in-plane lattice parameter of bulk MnO. The frequency
modulation atomic force microscopy (FM-AFM) image [Fig. 2.2(b)]
shows that the MnO(100) surface is atomically flat and consists of
terraces with lateral dimensions up to 500 Å
[3]. The MnO
stoichiometry of the oxide film has been confirmed by the
characteristic fingerprint provided by core level photoemission
spectroscopy (see Ref. [3] for details) and by the comparison
between the valence band photoemission spectrum and the
calculated [15] (PBE+Hubbard U method) density of states of MnO
given in Fig. 2.2(c), which provides further evidence of the good
quality of the MnO(100) multilayers. The main spectral features,
namely the triple peak structure with maxima at 2.2, 3.5 and 4.9 eV,
are reasonably well reproduced by the theory, although the
calculated peak with highest binding energy is upshifted by about 1
eV with respect to the experimental one. Finally, in panel (d) of Fig.
2.2 the HREELS phonon spectrum exhibits a main phonon loss peak
at
meV and a weaker structure at
meV, which can be
connected with the bulk optically allowed phonons at 33.6-36.4 meV
(Ref. [20]) and 62 meV (Ref. [37]). This finding is in line with the
single peak structure at 71 meV detected in MnO(100) crystals [49].
The comparison of the phonon spectrum of the MnO(100) thick
films with those of the MnO(100)-related monolayer structures
discussed
Mn O
previously
/Mn
O
[c(4 2)-Mn O
and
chevrons-
] clearly indicates that in the monolayer
regime the dipole-active phonon mode is strongly redshifted by
30
about 20 meV. A similar thickness dependence of the main phonon
frequency has been also reported for the MnO(100)/Pt(111) system
[38].
Fig. 9 30 ML MnO(100) on Pd(100) [3, 15]. (a) SPA-LEED pattern, recorded
with an electron energy of 90 eV. (b) FM-AFM image of the MnO(100) surface
(size 2500 2500 Å , detuning
= 5 Hz, bias U= 0.2 V). (c) Valence band
spectra excited by a photon energy of 120 eV (solid line) compared with the
calculated density of states (dashed line). (d) HREELS spectrum of the MnO(100)
film. Figure adapted from Ref. [3]
We have shown above that the number of oxidation states
available to manganese in the solid state is reflected in the complex
stoichiometric and architectural flexibility in the interface-dependent
monolayer regime. Similarly, the manganese oxide surface can
support a range of manganese oxidation states and oxide phases
because of the similar electronic character of the simplest Mn oxides
(MnO, Mn O and Mn O ), which ultimately differ by a diverse
electronic population and splitting of the partially filled
shell.
Langell et al. have shown that by proper adjustment of temperature
and oxidizing/reducing conditions it is possible to select the
dominant oxide forms, MnO, Mn O
or Mn O , present at the
MnO(100) surface [49]. Within this context, it is not unexpected that
under suitable preparation conditions (5 10
-2 10
mbar O
and elevated temperature) the Pd(100)-supported MnO(100)
31
structure can be converted into the Mn O (001) surface [16]. It is
nonetheless more surprising that either by annealing the MnO(100)
films at elevated temperature ( 770 K) in vacuo or by reactive
evaporation at relatively low pressure (
1 10
mbar) the
transformation of the MnO(100) structure into a MnO(111) surface
occurs [3]. The discussion of these two phase transitions will be
addressed in the following two sections.
10.2.2.1 Formation of Mn O on MnO(001)
Although, like in other spinel-like structured compounds,
terminations are formally considered to be unstable because of the
uncompensated electrostatic potential perpendicular to the surface, it
is reported that the cleavage of
leads distinctly to a (001)-
oriented surface [18]. Indeed, Noguera has rationalized that this
formal instability is an artificial consequence of the oversimplistic
ionic model and can be removed by a number of natural mechanisms
involving structural and electronic reconstructions [62]. The
epitaxial growth of
(001) has been achieved on MgO(001)
[35, 36] and circumstantial evidence indicates that
-like
surfaces can be obtained by mild oxidation of MnO(001) single
crystals [49]. In this section we discuss the formation of
(001) surfaces on Pd(100)-supported MnO(100) (denoted,
equivalently, as MnO(001)/Pd(100)) by high temperature oxidation
at intermediate and high molecular oxygen pressures.
32
Fig. 10 Comparison between the initial 20 ML thick
and the
phase prepared by oxidation of
film on Pd(100)
. (a) Mn 3 core
level spectra excited by a photon energy of 180 eV. The corresponding exchange
spin splitting is indicated. (b) Mn L
-edge X-ray absorption spectroscopy (XAS)
spectra recorded in total yield mode. (c) HREELS spectra. The energies of the
main energy losses are reported. (d) 2D SPA-LEED pattern of the 20 ML thick
film and (e) of the
phase. Both diffraction patterns
were recorded with an electron energy of 125 eV. (f) Corresponding 1D SPALEED profiles centered on the (00) reflex and recorded along the [110] substrate
direction at E =125 eV. Adapted from Ref. [16]
The Mn 3 core level spectra shown in Fig. 6(a) indicate
that exposing for 20 min Pd(100)-supported MnO(001) layers (20
ML) to molecular oxygen at a pressure of 2
mbar with the
sample kept at 770 K induces a marked decrease of the 3 splitting
down to 5.4 eV, a value that is in line with the measured exchange
splitting for bulk
(5.3-5.5 eV) and with the expected change
in the electronic configuration of the
evidence for the formation of
provided by the Mn
shell. Further experimental
-like films upon oxidation is
-edge X-ray absorption data reported in Fig.
6(b): as it is evident, the characteristic fingerprint of MnO (bottom
spectrum) is completely lost upon oxidation (top spectrum), and the
intensity ratio of the
Since the I(
)/I(
to
lines is reduced from
3.8 to
2.6.
) ratio can be related to the occupancy of the 3
33
orbitals on the Mn ion [48], the overall changes are again consistent
with the (partial) presence of Mn ions with 3
particular, the profile of the
configuration. In
peak is considerably broadened and
exhibits a triplet fine structure which is a distinctive signature of
[29]. The
stoichiometry of the oxidized film is also
confirmed by the HREELS phonon spectrum [Fig. 6(c)]. While the
MnO(001) film is characterized by a Fuchs-Kliewer phonon loss
peak at
65 meV and a much weaker structure at
48 meV (the
peak at 130 meV is a double loss of the 65 meV peak) as noted
above, the oxidized film is dominated by a new phonon loss at 83
meV, very close to the analogous vibrational peak at 650-654 cm
( 81 meV) that has been detected by Raman spectroscopy
measurements from mineral and synthetic hausmannite
crystals [17, 43].
The issue of the surface symmetry and growth direction of
the
layers is successfully addressed by SPA-LEED
measurements, see Fig. 6(d)-(e). Upon oxidation the sharp (1 1)
square pattern characteristic of well-ordered
evolves into
a weaker (2 2)-like pattern with broader spots, which preserves the
square symmetry of the unit cell but with a larger periodicity in the
real space. The line scan along the < 110 > high symmetry direction
[panel (f) of Fig.6] indicates that the spots appearing at roughly
positions are the first order spots of
53 1% of the
2
, which are located at
Brillouin zone and in real space describe
a square lattice with a unit cell parameter of 5.9 0.1 Å. From the
34
geometry of the bulk hausmannite spinel structure it is inferred that
the
film is (001)-oriented, with the
and
sides of the unit
cell aligned along the < 110 > directions of the MnO(001) substrate.
The data also demonstrate that due to the
the
9% lattice mismatch with
2 supercell of the underlying MnO(001) the supported
phase grows strained relative to the bulk phase (5.76 Å).
The DFT-based study of this system complements the
experimental information, giving additional insight into the atomiclevel characterization of the isolated
interface
surfaces and of the
[16].
In
line
experimental observations, the (001) orientation of
with
the
is found
to be the energetically most favorable, with a surface energy more
than 40 meV/Å
lower than that of the (110) and (100) surfaces.
Among the two possible as-cleaved (001) terminations, the mixed
oxygen and manganese surface with stoichiometry Mn O is more
stable than the manganese-terminated one by about 20 meV/Å .
The
structural
modifications
of
the
Mn O -terminated
crystal induced by the presence of the surface are
localized within the topmost two layers: they can be described in
terms of a buckling ( 0.2 Å) of the outermost species and of the
changes in the interlayer distances with respect to the bulk geometry
(in particular, there is a very pronounced contraction of 26% of the
first interlayer distance, partially compensated by an appreciable
expansion of 9% of the second interlayer distance) [16].
35
Fig. 11 Model structure for the interface between
and
and calculated relative stability (PBE). Black and red
circles indicate Mn and O atoms respectively. From Ref. [16]
Table 10.1: Calculated interlayer distances (d ) for the
optimized structure, expressed either in Å
(first row) or in percent with respect to the bulk interlayer spacing
of the strained Mn O
structure (second row). d
indicate the distances between the two
first two
d
and d
layers and the
layers at the interface, respectively, whereas
is the interfacial distance between the
and
contact planes.
1.26
2.27
0.74
0.95
0.95 1.02 1.00
-
-0.1
-0.1
29.5
-
+20.0
Based on the two distinct
1.10
-0.1
1.02 2.19
0.0
0.0
-
+0.1
terminations and
considering that the MnO(001) surface possesses a single
termination with planar unit containing an equal number of
36
manganese and oxygen atoms, two models of the interface emerge
naturally: (1) interface A, which is constructed by placing the Mnterminated (Mn-t)
surface on the
substrate
and (2) interface B, which directly joins the MnO surface layer with
the
-t layer of
. Different junctions can be
constructed by changing the registry of the interfacial
slice. The resulting models are sketched in Fig. 7: submodels A1 and
B1 correspond to an atop site registry of the
manganese atoms on the
interfacial
oxygen atoms underneath,
whereas submodels A2 and B2 refer to the corresponding bridge site
-
registry. The calculated relative stability of the
explored models, listed in Fig. 7, indicates that interface B1 has the
lowest energy and represents by far the most favorable model.
The
structural
characteristics
of
the
interface are summarized in Tab. 1
(optimized geometry) and Fig. 8 (layer by-layer decomposed charge
density plot) [16]. The structural optimization was performed using
a (2 2) unit cell relative to the MnO lattice constant and preserving
the bulk internal geometry in the three layers at the bottom of the
slab. A large vertical contraction of 20% of the
region of the
slab is obtained, which compensates for the in-plane tensile strain.
As a consequence of the latter strain, the distance between adjacent
surface oxygen rows increases by 0.7 Å, in contrast with the
contraction observed in the unstrained
surface.
37
Fig. 12 Charge density plots for the
interface: (a) layer-by-layer decomposition of the charge density
along the [001] direction. The dashed lines A, B, C and D (only A
and D labelings are given in the figure) are 2D projections of the
(001) planes displayed in panel (b) and (c) showing vertical cuts of
the charge density along the [100] and [010] directions,
respectively. Dark gray (red) and light blue circles indicate the
position of oxygen and manganese atom, respectively. The arrows
highlight the most significant internal structural relaxations. The
yellow line marks the separation between the
and
components, whereas the blue line indicates the unrelaxed starting
position of the topmost
layer. From Ref. [16]
In terms of interlayer distances it is found that the relaxations
are well localized within the first three layers and do not affect the
rest of the slab. By referring the changes of the interlayer distances
to the bulk interlayer spacing corresponding to the bulk Mn O
strained structure (i.e. with
), a significant compression of
the first interlayer distance (
= -29.5 %) is observed, which is
partially canceled out by the expansion of
down in the
(+20.0 %). Deep
part of the slab the situation remains practically
unchanged, including the distance
interface, the distance
between I2 and (S-7). At the
between the contact
layers is found to be slightly larger than the
and
interlayer
distance. The formation of the interface leads to an upward shift of
38
the manganese atoms placed in the MnO layer, thus inducing a
relatively small buckling of 0.2 Å between these atoms and the
interfacial MnO layer oxygens. It is worth noting that the bare
surface exhibits a larger displacement of the topmost
manganese atoms and an almost doubled surface buckling [15]. The
rest of the MnO slab remains unaffected and reproduces the bulklike behavior. In conclusion, the growth of strained
films on a
(001)-oriented substrate inhibits the structural
relaxations observed in the clean
surface. Finally,
because of the large vertical compression experienced in the
side of the slab there is a general in-plane
rearrangement of the atoms. As shown by the yellow arrows in Fig.
8 (b) and (c), the displacements mainly involve the movement of
oxygens along the [010] direction, which leads to an important
distortion of the octahedral O-Mn-O chain.
10.2.2.2 Epitaxial Stabilization of MnO(111) Overlayers
This conclusive section is devoted to the transformation of
MnO(100) layers into a polar MnO(111) surface, which is observed
upon annealing at high temperature in ultra high vacuum (UHV) [3,
4]. Keeping in mind the intrinsic thermodynamic stability of the
neutral MnO(100), this is an interesting and somewhat unexpected
result and therefore deserves an in-depth analysis. As already
mentioned, the problem with the stability of polar ionic crystal
39
surfaces has been long recognized, and recently its main lines have
been discussed in a number of review articles [62, 33, 14]. However,
while the (111) surfaces of rocksalt oxides such as MgO [31, 44, 7,
32], NiO [11, 12, 54, 13, 78], FeO [65, 68, 30], and CoO [52, 53]
have been intensively investigated experimentally and theoretically,
there is instead surprisingly little work done on the polar MnO(111)
surface. Using the grazing incidence X-ray scattering technique,
Renaud and Barbier have found that the MnO(111) single crystal
surface is non-stoichiometric, containing a mixture of MnO(111)
and Mn O (111) phases [67]. Rizzi et al. have grown epitaxial
MnO(111) films on a Pt(111) substrate and characterized their
structure by X-ray photoelectron diffraction (XPD) and LEED,
suggesting that the non-reconstructed MnO(111) surface can be
stabilized by adsorption of OH groups [69, 70]. Finally, DFT-based
thermodynamics has found that a (2 2) octopolar reconstructed
MnO(111) is the most stable surface over a wide range of the
oxygen chemical potential [25].
LEED patterns corresponding to manganese oxide films of various
thickness obtained by reactive evaporation of Mn at low oxygen
pressures and high temperature are displayed in Fig. 2.2.2. The
LEED reflexes now exhibit a pure hexagonal arrangement due to the
MnO(111) overlayer, and the presence of a faint streakiness signals
that faceting occurs to some extent. The phenomenon of the
formation of facets becomes more pronounced at higher oxide
coverage and high post-annealing temperatures, as visible in MnO
films of about 30 ML [Fig. 2.2.2 (b)-(d)]. A detailed analysis [3]
40
reveals that the facet rods are inclined with respect to the (00) rod of
the MnO(111) surface by an angle of 53
, which is very close
to the angle of 54.7 , expected between bulk (111) and (100) planes.
This clearly demonstrates that the facets are of (100) type.
Geometrical arguments, supported by FM-AFM profiles (not
shown), indicate that the MnO(111) surface is covered by triangular
pyramids exposing neutral (100) facets with the lowest surface
energy. Therefore, the formation of the (100)-faceted pyramids on
top of the MnO(111) surface provides a channel for minimizing the
total energy. Similar (100)-faceted pyramids have been observed
with AFM by Mocuta et al. on NiO(111) films supported on an Al O (0001) surface [54].
Fig. 13 (a) Conventional LEED picture of
Pd(100) evaporated at 670 K in
10 ML MnO(111) on
(O )=1 10
mbar and
annealed in UHV at 770 K. [(b)-(c)] SPA-LEED patterns of 30 ML
MnO(111) evaporated at 670 K in
(O )=1 10
mbar and
annealed at 820 K, recorded with electron energies of 54 eV and 70
eV, respectively. (d) SPA-LEED pattern of the surface in (b) and (c)
annealed at 920 K in UHV, recorded with an electron energy of 70
eV. Adapted from Ref. [3]
However, the replacement of the MnO(100) surface by the
polar MnO(111) cannot be completely explained by the faceting
41
argument alone. To understand this result we need to take into
account the strain energy at the metal-oxide interface, which is
determined by the lattice matching conditions. The MnO(100)Pd(100) interface is characterized by a large lattice mismatch of
14% along the [011] rows [Fig. 2.2.2(a)], whereas for the
MnO(111)-Pd(100) interface in one direction the lattice mismatch is
only
1%, which results in an almost perfect row matching along
the [01 ] rows, as illustrated in Fig. 2.2.2(b). This lowers the energy
of the MnO(111)-Pd(100) interface and presumably stabilizes the
formation of the MnO(111) overlayer. This finding is corroborated
by the DFT results of Fig. 2.2.2(c) showing the pressure dependent
variation of the energy of formation E
for MnO (100) and (111)
thin slabs. Although the (100) termination remains the most stable
surface in the whole range of variation of the lattice constant , the
curves show that there is a clear tendency of MnO(111) to gain
stability with respect to MnO(100) layers upon compression, i.e.
when the MnO bulk lattice constant shrinks in order to match that of
Pd(100) the energy difference E = E (111) - E (100) decreases
by a factor of 2.5 with respect to that calculated for
. This
result shows that when considering the stabilization mechanism of
polar surfaces of oxide films, the metal-oxide interface can play a
crucial role. Indeed we have demonstrated above that well-defined
and well-ordered MnO(111)-like interfaces are formed at monolayer
coverage, which can provide structurally graded interfaces for the
growth of thicker MnO(111) polar films. Importantly, the epitaxial
42
stabilization mechanism may also allow the design of oxide surface
orientations in thin films that are not stable in the bulk form. These
oxide phases can have interesting implications in the application of
ultrathin
oxide
films
in
diverse
areas
of
the
emerging
nanotechnologies.
Fig. 14 Real-space model of the (a) MnO(100)-Pd(100) and (b)
MnO(111)-Pd(100) interfaces. Note the row-matching condition
along the [011] rows at the MnO(111)-Pd(100) interface. DFT+U
derived plot of the formation energy of freestanding 3ML MnO(100)
and MnO(111) layers as a function of the lattice constant. From Ref.
[3]
Acknowledgements
We are very grateful to all coworkers mentioned in the references, in
particular to F. Netzer, S. Surnev, R. Podloucky and G. Kresse for
their scientific vision, the unfailing inspiration and the enthusiastic
support, and to F. Li, G. Parteder and V. Bayer for their invaluable
assistance during different stages of this work. Financial support by
the Austrian Science Funds FWF, by the 6th Framework
Programme of the European Community (GSOMEN and ATHENA),
and by the 7th Framework Programme of the European Community
(ERC Advanced Grant SEPON) is thankfully acknowledged.
43
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