Article
pubs.acs.org/cm
Interfacial Charge Transfer and Chemical Bonding in a Ni−LaNbO4
Cermet for Proton-Conducting Solid-Oxide Fuel Cell Anodes
Despoina-Maria Kepaptsoglou,*,†,∥ Kianoosh Hadidi,† Ole-Martin Løvvik,†,‡ Anna Magraso,§
Truls Norby,§ Anette E. Gunnæs,† Arne Olsen,† and Quentin M. Ramasse*,∥
†
Department of Physics, University of Oslo, FERMiO, Gaustadalleen 21, NO-0349 Oslo, Norway
SINTEF Materials and Chemistry, Forskningsveien 1, P.O. BOX. 124 Blindern, NO-0314 Oslo, Norway
§
Department of Chemistry, University of Oslo, FERMiO, Gaustadalleen 21, NO-0349 Oslo, Norway
∥
SuperSTEM Laboratory, STFC Daresbury Campus, Keckwick Lane, WA4 4AD Daresbury, U.K.
‡
S Supporting Information
*
ABSTRACT: In this work, we present an atomic scale study of the
structural and chemical characteristics of interfaces between Ni and
LaNbO4 grains in Ni−LaNbO4 cermets, a model composite material for
anodes in proton-conducting solid-oxide fuel cells (SOFC). Electron
energy loss spectroscopy (EELS) performed in an aberration-corrected
scanning transmission electron microscope reveals the absence of
reaction or interdiffusion layers at the interface. Changes in the valence
state of Ni as well as in the electronic structure of La, reflected by
changes in the EELS fine features at the interface, are shown to be
related to charge transfer across the interface. The experimental results
are in excellent agreement with ab initio calculations based on density
functional theory, which predict that direct chemical bonds are formed between the metal and the ceramic at this abrupt
interface, resulting in a redistribution of electronic charge across the interface.
KEYWORDS: scanning transmission electron microscopy, electron energy loss spectroscopy, density functional theory, charge transfer,
proton-conducting solid oxide fuel cells, electrodes, metal/oxide interfaces, LaNbO4
1. INTRODUCTION
The demand for efficient and clean energy technologies has led
to an increased interest in the direct conversion of gaseous fuels
into electricity using fuel cells. High-temperature solid-oxide
fuel cells (SOFCs) can be divided into two subcategories,
oxygen- and proton-conducting SOFCs, depending on the type
of conductivity of the electrolyte material of the cell (oxidic or
protonic, respectively).1 Proton-conducting solid-oxide fuel
cells (PC-SOFCs) employ acceptor-doped oxides with a high
density of oxygen vacancies, where protons dissolve as
hydroxide defects in the oxide at the expense of the vacancies.2
The oxide then becomes proton-conducting as the protons are
able to jump between their oxide ion hosts. In PC-SOFCs, the
electrical current is produced by the electrochemical oxidation
of H2 fuel to protons in the anode, which typically consists of
porous composite materials,3 such as cermets (ceramic−metal
composites). It is thought that the H2 is dissociatively adsorbed
from the fuel gas phase onto the metal or ceramic surfaces at or
near confined spatial sites, called three-phase boundaries, where
electrolyte, gas, and electrode are in contact.3−5 The porosity of
cermet anodes provides for this purpose plenty of three-phaseboundary contacts between the percolating phases of the
ceramic proton conductor, the metallic electron conductor and
the gas phase. Although the effects of geometry and kinetics of
fuel cell electrodes are widely studied,3,6−10 the role of the local
© 2012 American Chemical Society
electronic structure of metal/oxide interfaces and three-phase
boundaries in determining the cermet anode functionality is not
yet fully understood.
In particular, the interface between a metal and an oxide is
often characterized by a redistribution of electronic charges
across the interface. The spatial extent of this phenomenon
depends on the specific system and can arise from various
causes such as chemical bonding, formation of secondary
phases, or differences in the material work function on either
side of the interface. The description and understanding of this
charge transfer (which can be positive or negative), while
essential for the design and further optimization of PC-SOFCs,
is rather complex: different physical and chemical mechanisms
can be at play. To a first approximation, the contact can be
described electronically in terms of a Fermi level alignment
between the oxide and the metal.11 This alignment results in
the formation of a space charge layer at the oxide surface and
depends on the metal work function and the band gap of the
oxide. Locally, the charge redistribution is governed by the
presence of discrete energy states in the oxide band gap, which
originate from delocalized metal wave functions. These energy
Received: July 23, 2012
Revised: October 1, 2012
Published: October 1, 2012
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states are referred to as metal-induced gap states (MIGS)12 and
normally extend only a few atomic layers across the interfacial
region. Further parameters to be taken into consideration are
chemical bonding as well as local structural changes (such as
lattice distortions), as all will result in the creation of local
electronic interfacial states that could in turn affect the transfer
of charges.
In this work, we make use of the subangstrom spatial
resolution of analytical aberration-corrected scanning transmission electron microscopy (STEM) in order to probe the
structure and electronic configuration of a metal/oxide interface
at the atomic scale. The material selected for this study is a
cermet composite combining Ni- and Ca-doped (0.5% at.)
LaNbO4 (hereafter referred to as LNO for simplicity), a
potential electrolyte in PC-SOFCs.2,13−15 LNO exhibits
relatively high and pure proton conductivity, which, in
combination with its excellent chemical stability and chemical
compatibility with Ni,16−18 make it a good case study material
for the investigation of the electronic structure of interfaces in
PC-SOFC anodes. We combine state-of-the-art STEM imaging
and electron energy loss spectroscopy (EELS) with ab initio
calculations of the electronic and atomistic structure of a Ni/
LNO interface model, in an effort to shed light on the
interfacial bonding and charge transfer mechanisms at such
technologically important interfaces.
was subtracted using a 2-step baseline function, the slope of which was
determined by a linear fit across a 50 eV window at the continuum
after the L2 Ni white line. The 2-step function was then formed using
the continuum slope and steps with a 2:1 height ratio according to the
ratio of initial states that should result in a 2:1 ratio transition from the
2p3/2 and 2p1/2 to the continuum.24 The onsets of the steps were set at
the maxima of the L3 and L2 peaks, respectively. The ratio of the Ni
white lines was then determined by the ratio of the integrated intensity
of the baseline-subtracted signal, using Lorenz functions to fit the Ni
L3 and L2 edges.
The same fitting procedure was applied to spectra of pure Ni as well
as those with overlapping Ni and La edges as a validation of the
procedure. The edge overlap was separated by fitting the Ni L3 and La
M4 edges using Lorentz functions. This background subtraction
procedure was followed for all acquired spectrum images for which the
overlap in the spectrum images did not exceed 1−1.5 nm (as this
corresponds to the estimated delocalization tails of the La M4,5 signal
in our experimental conditions25) for individual spectra from the
linescans, as well as by averaging 2 to 3 adjacent spectra from each
linescan to improve signal-to-noise in the resulting processed data. The
experimental O K EEL spectra were compared with spectra calculated
with the FEFF9 code that uses the real-space full multiple-scattering
theory within the self-consistent muffin-tin potential approximation26
to simulate the near-edge region (ELNES). The spectra were
simulated using the initial state approximation (no core-hole) for
clusters including 150−200 atoms.
The atomic structure and density of states calculations were carried
out using spin-polarized density functional theory (DFT)27,28 through
the projected augmented wave method (PAW),29,30 as implemented in
the Vienna ab initio simulation package (VASP). The generalized
gradient approximation within the Perdew−Burke−Ernzerhof scheme
(PBE-GGA)31,32 was utilized to account for the exchange-correlation
corrections. Convergence tests confirmed that the numerical
uncertainty in the calculated total energy for the bulk unit cell of
undoped monoclinic LaNbO4 (LNOm) was less than l meV, when
using a 6 × 4 × 6 Monkhorst−Pack k-point mesh and a cutoff energy
of 850 eV. An undoped LaNbO4 model was used here for simplicity.
The addition of a small amount of Ca to LaNbO4 has been shown to
improve the conduction properties of the material, and at moderate
doping levels, Ca was found to segregate at the LNO domain
boundaries.2 However, no Ca was detected through EELS in the anode
material investigated for this study, likely due to the very small doping
levels involved. The electronic structure of the pure compound can
thus be considered a very good approximation to that of the
experimentally studied Ca-doped LNO. To reach the same accuracy in
the case of the Ni bulk unit cell, a Γ-centered 8 × 8 × 8 k-point mesh
and a cutoff energy of 370 eV were sufficient. The ionic geometries in
bulk Ni and LNOm were relaxed until the residual forces were less than
0.05 eV/Å. The ionic relaxation of the interface structure was
performed using only the Γ-point in k space and a convergence
criterion of 0.1 eV/Å for the remaining forces. For the final electronic
structure calculations of the total energy of the supercell, we used a
cutoff energy of 850 eV and a 3 × 3 × 2 k-point grid. Experimental
covalent radii of 1.69, 1.37, 0.73, and 1.21 Å for La, Nb, O, and Ni,
respectively, were used to project the local density of states (LDOS).
Soft PAW potentials were utilized,31,32 treating only the outermost
shells as valence electrons. The valence electronic levels were
5s25p66s25d1 for La, 4p65s24d3 for Nb, 2s22p4 for O, and 4s23d8 for Ni.
In order to model the Ni/LNOm interface structure, the
LNOm(101) and Ni(100) surfaces were merged, reflecting the present
experimental observations of sharp interfaces devoid of any
intergranular film. The model structure included 48 atoms in 4 layers
of LNOm and 75 Ni atoms distributed in 5 layers. The number of
layers for LNO was chosen so that the resulting LNOm slab would be
large enough to describe fully the (101) LNO surface. The reasoning
behind this choice is described in more detail in ref 4. The Ni lattice
was adjusted to the bulk relaxed lattice vectors of LNOm based on the
coherent method.33 The Ni primitive translation vector was thus
stretched along the longer lattice vector of the (101) LNOm surface at
the interfacial plane and compressed along the shorter direction. This
2. EXPERIMENTAL AND COMPUTATIONAL DETAILS
The anode cermet was prepared in pellet form by cold-pressing of NiO
and LNO powders in a 1:1 ratio and by sintering the resulting pellets
at a temperature of 1250 °C. The NiO−LNO pellets were
subsequently reduced to a Ni−LNO cermet by heat treatment at
700 °C in wet H2 atmosphere. Further details about the anode
preparation and performance are discussed elsewhere by Magraso et
al.14,16 Specimens for electron microscopy were prepared from the
reduced Ni−LNO pellets; 3 mm disks were drilled out of the center of
the pellets using an ultrasonic drill. The disks prepared by mechanical
polishing and ion beam thinning at 4 keV using a Gatan PIPS ion mill.
High-resolution high-angle annular dark field (HAADF) STEM images
and EELS were acquired from three instruments; TEAM 0.5,19 an
optimized double-corrected FEI Titan3 microscope equipped with a
Gatan Image Filter Tridiem, operated at a primary beam energy of 300
keV; a Cs-corrected VG HB501 dedicated STEM instrument20
operated at 100 keV equipped with a Gatan Enfina spectrometer;
and, finally, a Nion UltraSTEMTM100 dedicated STEM operated at
100 keV and also equipped with a Gatan Enfina spectrometer. The
inner and outer radii of the HAADF detectors in the conditions used
for the experiments were calibrated at 62−280 mrad for TEAM 0.5,
70−210 mrad for the VG HB501, 85−195 mrad for the UltraSTEM,
with probe convergence semiangles of 29, 20, and 31 mrad,
respectively. For the core-loss EELS, the collection semiangles were
32, 18, and 34 mrad, respectively.
In order to determine the ratio of the L3/L2 EELS edges of Ni at the
interface between Ni and LaNbO4 grains, the overlapping La and Ni
edges must be separated. Except where noted otherwise, the raw coreloss data were denoised using principal component analysis as
implemented in the HREM Research MSA plug-in for Digital
Micrograph.21 The spectra recorded on the VG HB501 instrument
were initially corrected for multiple inelastic scattering by the Fourier
ratio method,22 using low-loss EEL spectra recorded in identical
experimental conditions: the shorter inelastic mean free path at the
lower beam energy of this instrument made this additional step
necessary for accurate data interpretation. The background intensity of
the spectrum images was subtracted using a power-law baseline, and
the intensity of the spectrum image was normalized to the intensity of
the continuum after the Ni L2 edge. The methodology described by
Pearson et al.23,24 was employed for the determination of the Ni L3/L2
ratio. The remaining background intensity under the Ni white lines
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Figure 1. (a) HAADF STEM images, acquired in the TEAM 0.5 instrument, of a Ni/LNO interface; (b) close-up HAADF image of the same
interface showing the rotation of the LNO surface from (−101)m to (101)m across the twinned domains.
Figure 2. (a) HAADF STEM image of the Ni/LNOm interface and the path of the EELS linescan; (b) raw EELS spectrum image of the La M4,5 and
Ni L2,3 core loss edges, showing the La M5 and Ni L3 overlap; (c) Ni and La white line ratio values plotted as a function of position, together with the
intensity of the HAADF image along the linescan. The presented experimental data were acquired in the TEAM 0.5 instrument.
leads to lattice mismatches of −6.2% (compression) and 6%,
respectively, along these two lattice directions. The Ni lattice constant
was kept unchanged perpendicularly to the interfacial plane, only
adjusted by the monoclinic angle of LNOm in this direction. The
interfacial energy was calculated as 0.045 J/m2; such a low value of the
interfacial energy indicates that the supercell model is relatively stable
and thus physically relevant. More details about DFT modeling work
on the Ni−LaNbO4 system are provided by Hadidi et al.4
the t* subscript: at* = ct* = 5.40 Å, bt* = 11.66 Å, and
[010]m∥[010]t*. The transformation from the tetragonal to the
monoclinic structure upon cooling past 520 °C is accompanied
by the formation of a large density of twin-like domains parallel
to the (2 0 5.10)m/(5.10 2 0)m planes.34 Their formation has a
visible impact on the orientation of the LNOm surfaces and
interfaces; Figure 1a,b shows HAADF STEM images of a Ni/
LNOm interface. In such Z-contrast images, the intensity of
each atomic column can be directly related to the nth power (n
= 1.5−2) of its average atomic number Z and hence be used as
a guide for chemical identification.35 While the La−Nb bright
(heavy) columns are clearly resolved, the O positions could not
be observed in the HAADF images due to insufficient contrast.
All examined interfaces are faceted. The faceting is concave
and convex within the same grain in a zigzag manner in order to
accommodate the grain distortion due to the rotation between
adjacent domains. From the HAADF image (Figure 1b), it is
clear that the surface of the LNOm grain rotates sequentially
from one type of surface to another across each boundary (in
3. RESULTS AND DISCUSSION
At room temperature, LNO has a monoclinic Fergusonite type
crystal structure (space group C2/c, am = 5.56 Å, bm = 11.529 Å,
cm = 5.206 Å, αm = γm = 90°, and βm = 94.09°). Above 520 °C, it
undergoes a structural transformation to the tetragonal
Scheelite-type structure (space group I41/α, at = 5.40 Å, ct =
11.66 Å). For simplicity, crystallographic planes and directions
are henceforth denoted by m and t subscripts for the
monoclinic and tetragonal LNO phases, respectively. Furthermore, the tetragonal structure is hereafter described using
the axis of the monoclinic structure and will be denoted using
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Figure 3. (a) HAADF STEM image of the Ni/LNOm interface showing the EELS line scan trace; (b) raw O K core-loss spectra corresponding to the
linescan in panel a; (c) experimental EELS O K edge from the Ni/LNOm interface region and simulated O K spectra for several La, Ni, and O
containing compounds, carried out using the FEFF9 code. The presented data were acquired in the TEAM 0.5 instrument.
reported in the literature for metallic Ni.37−39 Approaching the
interface, the L3/L2 ratio increased up to 4 ± 0.3, a value that
agrees with those reported for NiO,37,39 suggesting a valence
change at the interface and the presence of Ni in a higher
valence state (although not necessarily a +2 state) in the
vicinity of the LNOm grain. The same increasing trend was
observed in all processed spectrum images recorded on all three
microscopes, in grains of different orientations, with absolute
values showing an overall deviation of 0.3. Moreover, a slight
shift of ∼0.5 eV toward lower energies can be observed in the
Ni L2,3 edge from the bulk of the Ni grain to the interface (inset
in Figure 2c). This shift is indicative of a change in the Fermi
energy of Ni: it is systematically observed (although smaller
than the energy spread of the instrument used to acquire the
data in Figure 2, 0.8 eV), including in data acquired on the VG
HB501 instrument (not shown) whose energy resolution is
0.35 eV. These observations thus add strength to the
conclusion that a change of valence state of Ni is taking place
at the interface.22
Variations in the Ni white line ratio could be explained by the
presence of an interfacial reaction layer such as NiOx, which
would indeed result in a local valence change. However, no
intergranular layer was directly visible in the HAADF images,
although as mentioned earlier, due to the polycrystalline nature
of the sample, few of the investigated interfaces were perfectly
abrupt. Furthermore, the fine features of the O K core-loss
EELS edge are known to be sensitive to the local bonding
environment. The presence of an oxide phase would therefore
be reflected in noticeable EELS signal variations. Figure 3a,b
shows an EELS linescan for the O K edge across a typical Ni/
LNO interface and clearly demonstrates that no oxygen is
detected on the Ni side of the interface. (See also Figure S1 of
Supporting Information.) Background-subtracted spectra from
the linescan (Figure 3b) reveal a drop in the total O intensity,
but no fine structure change as the probe approaches the
interface. In some other data sets (not shown) featuring a slight
Ni/LNOm grain overlap, a weak O K edge appears within the
overlapping region, but again without any significant change of
the edge fine structure.
In order to rule out the presence of an intergranular film,
FEFF9 simulations of the O K edge were carried out using the
code for LaNbO4 and a variety of possible secondary phases,
such as La2O3 and La2NiO4 (Figure 3c). The calculations were
performed for different crystallographic orientations, showing
Figure 1b from (101)m to (101)m and so on), following the
rotation of the domains. These surfaces correspond to the
(101)t* surface of the high temperature Scheelite phase and will
bend in an accordion manner during the inevitable thermal
cycling from preparation conditions via room temperature to
the eventual operating conditions of the fuel cell. The rotation
of the surfaces can thus be expected to introduce further strain
and defects at the metal/oxide interface. Although no
significant chemical bonding differences are expected at the
metal/oxide interface between pristine tetragonal and monoclinic phases,5 elastic strain is expected to influence ionic
mobility across the heterointerfaces.36
The electronic structure of the Ni/LNOm interfaces was
assessed experimentally by STEM EELS. The white line
intensity ratio of transition metals is known to be sensitive to
their valence state.21 Therefore, the determination of the L3/L2
intensity ratio for Ni can provide valuable information about its
oxidation state at the interface. Core-loss spectra were recorded
along lines across Ni/LNOm interfaces; many such data sets
were acquired with the LNOm grains in different crystallographic orientations (here showing [010]) in order to exclude
orientation and channelling effects, while post mortem images
were recorded after the EELS acquisition in order to exclude
beam-induced damage to the sample. As the only Nb EELS
edge (M4,5) within the usable energy range for these
experiments is mostly devoid of any fine features whose
changes could indicate electronic reconfiguration, the rest of
this discussion will focus on the La, Ni, and O components of
the metal/oxide interface.
Some as-recorded La M4,5 and Ni L2,3 EEL spectrum images
(an example of a linescan is presented in Figure 2a,b, and inset
in c), showed a limited overlap of the La M4 and Ni L3 intensity
around the interface. As no reaction layer could be directly
observed from the HAADF images, the core-loss edge intensity
overlap is attributed to grain overlap. The polycrystalline nature
of the cermet sample and the random orientation of the grains
made the detection and imaging of a perfectly abrupt interface
challenging. Some degree of overlap between the La M4 and Ni
L3 edges is also expected due to the large EELS tails for the La
M4,5 edge signal.25 Figure 2c shows the plot of the Ni L3/L2 and
La M5/M4 ratios corresponding to the linescan in Figure 2a, as
well as the simultaneously acquired HAADF intensity along the
linescan. The L3/L2 ratio deep within the Ni grain was
determined to be 3.3 ± 0.2, in good agreement with values
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Figure 4. Ball and stick models showing the simulated interface between the (101) LNOm surface and (100) Ni surface (a) unrelaxed and (b)
relaxed.
conflicting interpretations.47 However, DFT calculations can be
used in order to rationalize the electronic structure variations
observed experimentally. The experimental HAADF images
thus motivated the creation of an atomistic LNOm(101)/
Ni(100) interface model by merging directly the LNOm(101)
and Ni(100) surfaces (Figure 4a,b). Although the supercell
model was complex, including 48 and 75 atoms in 4 and 5
layers of LNOm and Ni, respectively, the final relaxed structure
was found to possess an interfacial energy of only 0.045 J/m2.
Such a low value of the interfacial energy, coupled with the
minimal distortions of the Ni lattice, demonstrate the relative
stability of this supercell and confirm the physical plausibility of
the model. Table 1 presents the calculated interatomic
no significant difference with orientation; the simulated spectra
plotted in Figure 3c are therefore averaged spectra calculated at
the so-called magic angle, with no additional broadening
applied. The simulated spectra in Figure 3c are plotted together
with an experimental O K spectra from the Ni/LNOm interface
(second from bottom curve in Figure 3b), bulk LNOm from the
same experiment and NiO from a standard specimen. None of
the simulated edges for the secondary oxide compounds exhibit
the broad overall shape of the experimental O K edge, which in
turn fits rather well the simulated fine structure for the O K in
LNOm. Futhermore, the experimental O K edges at the
interface are distinctly different from those acquired from the
NiO specimen: the latter presents a rather large intensity hump
at ∼543 eV (see also ref 40), which is clearly not present in our
experimental results. From the above, it becomes clear that the
apparent oxidation change of Ni is unlikely to be a consequence
of the presence of a NiOx interfacial layer: another mechanism
must be at play.
Another interesting observation arising from the EELS study
is that the white line ratio of La shows a decreasing trend
toward the metal/oxide interface, from a value of 1.0 well inside
the LNOm grain. In rare earth elements, the M4,5 lines
correspond to transitions from the 3d orbitals to unoccupied
4f states (3d104fn → 3d94fn+1 transitions), which are highly
localized.41 In some cases such as Ce and Pr, this localization is
known to generate changes in the white line intensity ratio
depending on the occupancy of the 4f state.42−44 By contrast,
for La, which in general is found in a +3 oxidation state, no
variation of the La white line ratio has ever been reported, to
our knowledge, even though La2+ compounds are known to
exist.45 In particular, Mikheeva et al. suggest that no significant
change should occur in the 4f band during La oxidation and
that occupancy changes should only be expected in the 5d
states, therefore not affecting the M4,5 white line intensity ratio
in EELS.46 Nevertheless, this La white line ratio change was
consistently observed at all interfaces including those with no
significant grain overlap, and it must therefore reflect a change
in the electronic structure of La, although perhaps not a pure
valence change. EELS simulations for the La edges (and more
generally for rare earth elements) are notoriously difficult to
match accurately to experimental spectra and can lead to
Table 1. Interatomic Distances (in Å) between Ni and LNO
Elements within the First Coordination Shell, for the
Relaxed Interface Structure, Compared to Experimental
Values for Bulk Compounds
atoms
relaxed
experimental
Ni−La
2.85, 2.94, 2.98
2.92 in all La−Ni
compositions
Ni−Nb
2.7−2.98
2.69−2.79
Ni6Nb6O
Ni−O
1.86−1.98
2.08−2.95
La(NiO3), NiO,
NiO2
distances of the modeled LNOm(101)/Ni(100) interface,
compared to relevant experimental bond lengths for known
compounds.48 In the interfacial model, the calculated
interatomic distances between Ni and all constituting elements
in LNO were similar to, or smaller than, these experimental
interatomic distances. This is a strong indication that chemical
bonds are established at the interface between Ni−O, as
frequently observed in metal oxide heterointerfaces,49,50 but
rather interestingly also between La−Ni.
This interpretation is confirmed by the calculated density of
states (DOS) for the model interface. Figure 5 presents the
LDOS projected onto the La and O atoms at the interface layer
as well as that of their nearest Ni neighbors. Only the valence
orbitals reflecting chemical reactions have been plotted. The
valence orbitals of Ni clearly overlap with the valence orbitals of
La (Figure 5b,e) and O (Figure 5a,d). These overlaps can be
considered as indirect evidence of chemical bonding between
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Figure 5. Local density of states projected on the atoms at the interfacial LNO and Ni layers (panels, a−e). Each panel presents a specific angular
momentum orbital of an interfacial La (solid black) or O (dashed red) atom along with its nearest neighbor Ni (solid magenta) atom. Panels c and f
compare the La electronic structure at the interface and in the bulk. The La d and f orbitals are shown for the interfacial layer (La-1; black), the
second LNO layer from the interface plane (La-2; gray), and for La in bulk LNO (La-b; dashed green).
Figure 6. Contour plot of the electron localization function. A value of 1 means very high localization (low kinetic energy of the electrons), while 0
means no localization. The plots are drawn within planes containing intersecting neighbor Ni and La atoms (left) or neighbor Ni and O atoms
(right).
The changes in the white line ratio of La are arguably more
difficult to relate to the DFT results. Nevertheless, it is clear
from the calculated LDOS (Figure 5c,f) that the d and f orbitals
of La, both of which have unoccupied electronic states and
therefore contribute to the formation of the EELS signal,
undergo a considerable change when going from bulk LNOm to
the interface structure. For instance, comparing the La d
orbitals in the bulk and at the interface reveals that the LNOm
band gap disappears at the interface. This phenomenon is
Ni−La and Ni−O. It was recently shown that overlap
population (or crystal orbital overlap population) diagrams
can be used very effectively to qualitatively interpret X-ray
absorption near-edge spectroscopy (XANES) or EELS
spectra.51 The calculated overlap between the Ni and the La
or O atoms LDOS at the LNOm at the interface is thus likely to
be the cause of the change in valence of Ni, as experimentally
reflected by the increase of the EELS white line ratio of Ni.
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LNOm and is expected to be confined to one atomic layer at the
interfacial plane. Experimentally, the change in Ni valence and
the La electronic reconfiguration can be considered as direct
experimental indicators of this charge exchange. However,
because of the complexity of the electronic configuration at the
Ni/LNOm interface that cannot be fully described by a fully
ionic model, a direct quantitative correlation of the measured
white line ratio with the actual charge transfer is difficult to
perform. Experimentally, the charge transfer appears to extend
over several atomic layers (a distance of about 2 nm; Figure
2c). This extended charge transfer area can likely be attributed
to the variation of the relative orientation of the experimental
interface with respect to the electron beam throughout the
thickness of the TEM specimen. The interface is not perfectly
abrupt resulting in partial overlaps of the metal and oxide along
the beam direction, at least in some parts. The spatial
localization of the valence measurement might also be affected
by electron channelling or other delocalization effects of the
EELS signal and will in turn result in observing the charge
transfer further away from the interface.
This study shows that, given a stable coexistent pair of
electrolyte and electrode, it is possible to obtain directly a
thorough structural, compositional, and electrical characterization of electrode interfaces at an atomic level using advanced
analytical microscopy. Although this study does not specifically
account for the presence of protons, when combined with
atomistic and electron structure modeling, this approach opens
up unprecedented possibilities for understanding and interpreting the electrochemical characteristics and performance of
SOFC and PC-SOFC electrodes since the electronic structure
is decisive for the local solubility and diffusivity of protons
across the interface (see relevant work on hydrogen energetics
in ref 5). We expect that such studies will eventually be able to
predict activation barriers for rate-limiting steps in electrochemical redox reactions, for instance, to be compared with
polarization and impedance spectroscopy data from experimental investigations of the electrodes and devices under
operating conditions.
expected at a metal/oxide interface: the metal and the oxide
electronic structures undergo Fermi level alignment12 resulting
in the formation of MIGS,5,13 which, in this case, are projected
at the La d orbital. Moreover, in the conduction band, we
observe a 2.0 eV shift toward lower energies in the f orbital of
the interfacial La compared to that in the bulk (Figure 5f).
These changes in the LDOS reflect a strong rearrangement in
the electronic configuration of La close to the interface. Hence,
it is possible to assume that, in a complete electronic picture of
the La atoms at the interface, the rearrangement of orbitals, due
to the interfacial bonding, alters altogether the availability of the
final f states, and this could, in turn, be the cause for the drop in
La white line ratio observed in the EELS data.
We have thus clearly established the presence of direct
chemical bonds between the Ni and LNOm parts of the cermet
anode. In order to clarify the nature of the bonds, we have
calculated the electron localization function (ELF) at the
interface, which gives a direct spatial view of the bond
location.52 Contour plots of the ELF have been displayed in
Figure 6. The presence of electronic charge density between the
La and the Ni, as well as between the Ni and O atoms, is clear
evidence that a direct chemical bond is indeed established
between these elements at the interface. The high degree of
localization in the vicinity of the La and O ions rather than
around their Ni ion counterparts results from the presence of
delocalized valence electrons in the Ni d orbital. The ELF
results thus confirm that charge transfer is taking place from Ni
to LNO atom-wise, with direct chemical Ni−La and ionic Ni−
O bonds being responsible for the interface dipole. Hence,
Bader analysis was employed for the numerical estimation of
charge transfer between the atomic species at the interface.32
The sum of Bader charges for each layer relative to the bulk
(reference) state of Ni and LNOm is plotted in Figure 7, as a
function of position across the interface, the layers being
numbered separately for Ni and LNOm sequentially away from
the interface. This charge analysis reveals that a partial charge
transfer of 1.6 e for the entire simulated model (corresponding
roughly to 0.1 e per Ni atom) does take place from Ni to
4. CONCLUSIONS
In this work, the detailed atomic distribution and the electronic
structure at the interface between Ni and LaNbO4 has been
investigated, with particular focus on the origins of charge
transfer across the interface. This was achieved through a joint
experimental and modeling study, combining atomic resolution
scanning transmission electron microscopy and electron energy
loss spectroscopy investigations with density functional theory
calculations. The experimental results revealed a drastic change
of the electronic structure when moving from the bulk to the
interface region, quantified by a significant change of the white
line ratios of Ni and La, respectively. Despite the consistent
increase in the Ni oxidation state, the presence of a Ni oxide
layer was excluded. Thus, the interface was shown to be
chemically abrupt, leading to the assumption that bonding
occurs between Ni and the O of the LaNbO4 phase. The
experimental results were compared to DFT calculations of a
defect-free model of the Ni/LNO interface. They demonstrated
direct chemical bonding and a significant exchange of electrons
from Ni to La and O at the interface layer. This charge transfer
was quantified by theoretical calculations in very good
agreement with the experimental findings. The drastic changes
of the La f states at the Ni/LaNbO4 interface found in the
calculated LDOS are thought to be the cause of the change in
Figure 7. Calculated electronic charge transfer for the four layers of
LNO (red squares) and five layers of Ni (blue discs) in the interface
super cell. A Bader analysis was used to quantify the charge at different
atoms. A positive charge transfer means a gain of negative charge
compared to bulk atomic states. The layers are numbered sequentially
from the first interfacial layer to the next one along each component
(see also Figure 4). The charge differences between interfacial
neighbor layers have been indicated: lines are drawn as guide to the
eye only.
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Chemistry of Materials
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the La white ratio experimentally observed. These results show
for the first time how atomic-scale analytical electron
microscopy measurements coupled with ab initio calculations
can be used to characterize directly the nature of the chemical
bonds forming across a metal/oxide interface in a complex
cermet material. Such understanding is essential for further
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■
ASSOCIATED CONTENT
S Supporting Information
*
HAADF survey image and signal map; normalized O K edge
spectra; HAADF STEM images. This material is available free
of charge via the Internet at http://pubs.acs.org.
■
AUTHOR INFORMATION
Corresponding Author
*Tel: +44 (0) 1925 864 902. E-mail: dmkepap@superstem.org
(D.-M.K.); qmramasse@superstem.org (Q.M.R.).
Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS
This work is financed by the Research Council of Norway
under the NANOMAT program (project 182065/S10). Part of
this work was performed at National Center for Electron
Microscopy, Lawrence Berkeley National Laboratory, which is
supported by the Office of Science, Office of Basic Energy
Sciences of the U.S. Department of Energy under Contract No.
DE-AC02-05CH11231. The SuperSTEM Laboratory is funded
by the U.K. Engineering and Physical Sciences Research
Council (EPSRC).
■
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