Spectromicroscopy for Addressing the Surface and Electron Transport
Properties of Individual 1-D Nanostructures and Their Networks
Andrei Kolmakov1*, Sai Potluri1,
Alexei Barinov2, Tevfik O. Menteş2, Luca Gregoratti2, Miguel A. Niño 2, Andrea Locatelli2,
1 Southern Illinois University at Carbondale, USA; E-mail: email@example.com
2 Sincrotrone Trieste, 34012 Basovizza-Trieste, Italy
Keywords: SnO2 nanowires, nanobelts, surface, electron spectroscopy, microscopy, PEEM, LEEM, SPEM, LEED,
gas sensors, catalysis
Understanding size/dimensionality dependent phenomena and processes relevant to chemical sensing and
catalysis require analytical methods with high surface sensitivity, which can exploit the structure and
composition of nano-materials at their natural length scales and working conditions. In the present study
we explored the potentials and complementary capabilities of several surface sensitive microscopy
approaches to shed light on the properties of individual SnO2 nanowires and their networks. Our results
demonstrate the unique opportunities provided by synchrotron-based photoelectron microscopies for
surface sensitive structural and chemical analysis, including in-situ characterization of electron transport
properties of a nanostructure wired as an active element in chemiresistor devices.
Table of Content Figure
Implementation of “bottom-up” paradigm in technology is based on fabrication and assembly of
artificial objects with dimensions ranging from a few to a few hundreds nanometers with pre-designed
functionalities. Semiconducting quasi one dimensional (quasi-1D) nanostructures, in particular metal
oxide (MOx) nanowires and nanobelts, are among the most promising platforms for gas sensing, energy
harvesting/conversion and opto-electronic applications.1-3 When the effective diameter of these objects
shrinks to the nano-scale the surface starts to play a significant role and can even dominate their
properties. For example, the attractive novel properties of semiconducting metal oxide nanowires in gas
sensing4-8 and electronically tunable reactivity 9, 10 have been attributed to the intimate coupling between
surface redox processes and electron concentration inside the nanostructure. In order to control this
surface ↔ electron transport interplay and to develop a reproducible nanodevice the fundamental
relations between the surface composition, morphology and electronic structure of oxide nanostructures
along with their physical and chemical properties have to be identified. The complexity and challenges in
such studies are dictated by two major factors: (i) the purity (both surface and bulk) and (ii) the strong
size and morphology dependence of the observed phenomena. In addition, the substrate and electrodes in
the real world devices can induce significant changes to the electronic properties of the nano-structure.
Therefore, it is desirable to measure individual and pre-wired nanostructures, since the results obtained on
ensembles (bundles, films, etc.) lack size and morphology specificity.
Up to date the structural and compositional ex-situ and in-situ characterization of the oxide
nanostructures has exclusively been performed by means of a scanning electron microscopy (SEM) or
transmission electron microscopy (TEM) .11, 12 The imaging and analytical capabilities, combined with the
latest developments of sample stages in these machines have allowed in situ exploration of the
mechanical, compositional, electrical and other properties of the nanostructures, as well as high resolution
imaging and analysis at elevated pressures and temperatures. Both techniques provide very high spatial
resolution, but their spectroscopic abilities, based on energy filtered electron and photon emission, such as
Electron Energy Loss, Auger Electron and Energy dispersive x-ray spectroscopy, are limited in energy
resolution and/or surface sensitivity.
Alternative spectro-microscopy methods with surface sensitivity are those using x-ray sources
and detecting photoelectrons emitted from the top few layers of the specimen. In these x-ray
photoelectron microscopes information for the chemical state and electronic structure of the surface is
provided by the two well-developed spectroscopic methods: (i) X-ray Photoelectron Spectroscopy (XPS),
which uses monochromatic photons to obtain energy distribution curves of the photoelectrons emitted
from discrete atomic electron levels13 and (ii) X-ray Absorption Near-Edge Spectroscopy (XANES),
where the total electron yield spectra are measured while varying the photon energy around the x-ray
absorption thresholds corresponding to a given core electron binding energy14. Along with the surface
chemical sensitivity another advantage of these photoelectron-based spectroscopies, is that the radiation
damage, which can induce irreversible structural and chemical changes in the nano-object, is weaker
using x-rays relative to the electron beam. Probing the matter at sub-micrometer length scales using XPS
and/or XANES requires the intense and tunable x-rays provided at the third generation synchrotron
facilities. The spatial resolution is achieved using two different approaches, which classify the
microscopes as scanning and full-field imaging.15-17 In the scanning photoelectron microscope (SPEM)
the incident photon beam is demagnified to a submicrometer (ca 100-200 nm) spot onto the sample,
which can be raster scanned with respect to the microprobe. In the x-ray photoelectron emission
microscopes (XPEEM) a magnified image of the irradiated area is obtained by projection of the emitted
electrons using appropriate electron tailoring optics. These synchrotron-based microscopes can be
successfully complemented with conventional surface science methods for structural analysis, as low
energy electron diffraction (LEED) and low energy electron microscopy (LEEM). In fact, an important
advantage of XPEEM instruments is the possibility to perform LEEM in the same set-up18, which provide
complementary morphological information with spatial resolution of 10 nm and micro-spot LEED
monitoring the long-range atomic order from an area of 1-2 microns. SPEM and XPEEM have
successfully been used for analysis of C and MoSx nanotubes,19-25 GaN nanowire,26 oxide nano-particles
and wires.27, 28 A few experimental impediments for PEEM imaging and spectroscopy on an individual
pre-wired nanowire chemiresistor were reported in our previous work29 and some solutions were
This communication reports on the surface structure, composition and morphology of individual
quasi-1D SnO2 nanostructures, their percolating networks and corresponding transport properties when
they are indexed as a part of the device. The results, obtained by combining surface sensitive spectro-
microscopy and transport measurement methods, provide new important information for these nano-
materials. The surface and sensor properties of the macroscopic counterparts of the selected for this study
SnO2 nanostructures are already well known30 and we have developed fabrication and handling protocols,
which preserve their surface cleanliness (see Ref. 31 and supporting material therein).
Results and Discussion
The SPEM and X-PEEM/LEEM experiments were performed at the ESCAmicroscopy32 and
Nanospectroscopy beamlines17 at ELETTRA synchrotron radiation facility, hosting the SPEM and
XPEEM-LEEM microscopes, respectively. XPEEM-LEEM microscope is equipped with a hemispherical
electron analyzer, reaching an energy resolution of 200 meV, and provides the complementary
spectroscopic methods XANES and XPS. The lateral resolution is 10 nm in LEEM operation mode, and a
few tens nm in energy-filtered X-PEEM imaging. The instrument can also perform micro-LEED and
micro-XPS measurements, restricted to an area of 2 µm in diameter.
Surface Structure of the individual nanostructure
Figure 1(a, b) depict LEEM images of two different parts of the same SnO2 nanowirei. The panels (c) and
(d) represent the micro-LEED patterns collected at two locations (marked with white circles in (a) and
(b). Moving the sample in front of the LEEM objective, we were able to monitor the evolution of LEED
patterns along the nanostructure’s length. Both 1x1 micro-LEED patterns, marked by the white dashed
line, corresponds to a rutile surface unit cell with real space dimensions: (3.23≤0.1 Å) x (4.88≤0.12 Å) in
agreement with the ones from (010) or (100) facets of the SnO2 rutile crystals33. Based solely on the
i Here we are using the general term “nanowire” since the resolution is not sufficient to distinguish between
nanostructures with different morphologies
LEED pattern it is not possible to discriminate between these two facets. However, based on previous
HRTEM results showing that the preferable growth direction is  and (101) / (010) top/side faceting
of SnO2 rutile nanobelts and nanowires with rectangular cross section34, most probably the LEED pattern
corresponds to a (010) side facet of the nanowire. This assignment is supported by the facile bending of
the nanowire/nanobelt in this plane and the appearance of weaker diffraction spots in the LEED pattern.
The latter ones are an evidence of the single twin structure in the nanowire/nanobelt. Very similar SAED
patterns, taken along the  zone axis, have been previously reported34 for large diameter nanobelts
that agrees well with our assignment.
Another class of quasi-1D nanostructures, explored by LEED is nanowires, which change
their growth direction due to perturbed growth conditions, while preserving their
crystallographic orientation. In contrast to the nanobelt case, the V-shaped (“kinked”)
nanostructure in Fig. 1(e) shows the same LEED pattern independently of the position. This
observation corroborates nicely with previous HRTEM studies of the segmented and zigzag
SnO2 nanostructures which have shown that such structures are in fact “sculptured” single
crystals and their growth axis can be altered into crystallographically equivalent directions by
tiny changes in the growth conditions.35, 36 Again, the LEED pattern in Fig. 1(e) corresponds to
the (010) facet of the nanostructure. Assuming that the growth takes place along the [10-1] and
 directions, the angle at the kink is expected to be 67.9°, in agreement with the
experimentally determined value ~70°.
XANES and XPS on individual nanostructures
Figure 2 shows the XANES spectrum (green top line) of the SnO2 nanobelt, measured with X-
PEEM by collecting stacks of images while scanning the photon energy around the Sn 3d edge.
For comparison the XANES spectra of rutile SnO2 (red) and tetragonal SnO (blue) samples,
reported in Ref.37 are also shown. Analogous to the macroscopic counterparts, the XANES
spectrum from the individual nanostructure consists of the Sn 3d5/2 and 3d3/2 features (separated
by ~8.5 eV), giving rise to six major transitions into the unoccupied Sn electronic states. Close
inspection of the three spectra in Fig. 2 reveals that in spite of the general similarity there are
some differences in the relative intensity and shape of the nanobelt peaks, most evident in the 3-4
eV range. The enhanced relative intensity of the 4.2 eV peak, reported previously in Ref.37, has
been attributed to oxygen deficiency, induced by surface reconstruction. Since our LEED
patterns do not evidence a reconstruction, this suggests that the nanobelt surface is partially
reduced and the spectrum contains overlapping features of the SnO2 and SnO phases, as
indicated in the figure.
More detailed information for the surface chemical state has been obtained from the
photoelectron spectra of individual SnO2 nano-(meso-) structures, measured by SPEM, where the
incident photon beam is focused to a small spot (100-200 nm). SPEM can be operated in two
modes: (i) imaging by collecting photoelectrons within a selected kinetic energy window (Fig.3
a), while scanning the sample with respect to the focused beam and (ii) conventional energy
dispersive electron spectroscopy from the illuminated local spot (micro-XPS) (Fig.3 (b, c, d)).
When the element under consideration is present in a single chemical state the spatial variation in
the contrast of the images reflects the variation of the photoelectron yield, which is a measure of
the local concentration of the element. Fig.3 (a) shows the Sn 4d image of a SnO2 nanobelt
(bright) placed on a gold support (dark). Figure 3 (b) shows two representative sets of O 1s, Sn
3d5/2 and VB spectra measured on microspots on a freshly oxygen-plasma treated nanostructure
(bottom spectra) and after annealing to 5230K in UHV (top spectra). The spectra are
deconvoluted considering the reported binding energy of Sn 3d5/2 and O 1s core levels of SnO
and SnO2 macroscopic polycrystalline and single crystal samples.38-41 The presence of SnO-
derived components in the O 1s and Sn 3d5/2 spectra is in accordance with the XANES spectra in
Fig. 2. This should be attributed to lattice oxygen deficiency and reduced oxygen concentration
at the surface, becoming more prominent after annealing to 523K. The spectra from the fresh
oxygen-plasma treated sample have a weaker SnO component and contain also a third high BE
component, which dominates the O 1s spectra. According to the literature41 the rather broad
component at ~ 531.7 eV (marked with yellow in the Fig.3b) apparently originates from
chemisorbed O-containing species, such as Oδ-, O2δ- or OHδ-. Correspondingly, the high BE
component in the Sn 3d5/2 spectra can tentatively be attributed to surface Sn atoms bound to
adsorbed species, resulting in O coordination higher than in SnO2. The relatively high intensity
of the ‘Ochem’ component is due to the very high surface sensitivity, resulting from the low
kinetic energy of the O 1s and Sn 3d5/2 photoelectrons and the grazing acceptance geometry of
the energy analyzer in SPEM. The origin of this component is confirmed by its strong
attenuation due to desorption upon annealing to 523 K. The ‘Ochem’ component also gradually
losses intensity under prolonged x-ray irradiation at room temperature due to photon-induced
desorption, whereas the oxide components remain practically intact. The VB spectra in Fig. 3(d)
are measured with a lower surface sensitivity due to the high photoelectron kinetic energy. They
are dominated by the SnO2 main peaks centered at ~ 5, 8 and 11 eV, originating from the O(2p)
non-bonding, Sn(5s)-O(2p) and Sn(5p)-O(2p) states, respectively, whereas the SnO-related
features at ~3, 6.5 and 9.7 eV are appearing as weak shoulders.42 The only distinct difference in
the VB spectra of a fresh and annealed nanostructure is the increase of the emission at ~ 3 eV,
attributed to Sn(5s-5p) hybridized states in SnO band gap.37, 41 It is clearly manifested by the VB
spectral difference shown by the dashed line in panel (d). These results support that the oxygen
deficiency is mostly at the surface and due to the larger probing depth is not very evident in the
Addressing the electron transport in working device
When the SnO2 nanostructures are wired as active elements in nanodevices, their electronic
and surface properties can not be fully decoupled from effects induced by other parts of
nanodevice such as electrodes, support etc. Some examples of possible effects on the nanowire
properties are the electrostatically controlled reactivity,43 the gate induced electrodesorption10
and the influence of the mobile parasitic charges on the transport properties.44 In order to get
insight how these phenomena affect the electronic status of the nanostructures, wired as part of a
device, the analytical measurements have to be performed under conditions close or identical to
the working ones. The straightforward application of the photoelectron spectromicroscopy to
working nanodevices is generally hampered by parasitic charging, since the standard device
architectures routinely involve isolating layers or support in the proximity of the nanostructure.45
One of the possible solutions is to use a model UHV compatible device containing suspended
nanostructure similar to the one described in Ref.29. For the present study we fabricated a model
chemiresistor gas sensing device using Si3N4 membranes with periodic 2 microns holes, as a
support media for the nanowire network (Fig. 4a). The size of the holes is an order of magnitude
larger than the x-ray microprobe of SPEM, which allows the nanowires placed over the holes to
be locally probed. The SPEM experimental setup for imaging the SnO2 nanowire network
contacted by two electrodes is depicted in Fig. 4a. For this set-up the conductivity between the
electrodes has a percolating character and depends sensitively on the gas environments and on
the density and transport quality of the straight parts of the nanowires and their junctions
(nodes)44. Under such conditions the energy position and the width of the photoelectron spectra
can be affected by external bias, band bending and/or local electrostatic potential at the irradiated
point. For zero bias voltage, the magnitude of the steady state potential depends on the
conductivity and electron yield, which is a function of the photon flux, photo-ionization cross
section and secondary electron emission. For low conductive materials, since the screening of the
created electron holes is not sufficient, the electron emission results in local positive charging,
which results in reducing the kinetic energy of the emitted photoelectron with the same value.
Since the ‘disturbed’ sample volume by the x-ray induced electron emission is of the order of
~50 nm, apparently the screening becomes an issue with reducing the dimensions of the
semiconducting oxide structure. In fact, for the measurements described in previous sections the
whole nano-structure is placed on a grounded Au foil, the contact with which provides the
electrons for screening the created electron holes. Thus the Sn 3d5/2 and O 1s spectra in Fig. 2
and 3 appear at binding energies identical to that measured of the extended materials. For the
transport experiments only one side of the nanowire is in contact to the grounded gold electrode.
Due to the low conductivity and significant length of the nanowires the screening of the core
holes is not sufficient anymore and the corresponding Sn 3d spectrum in Fig. 4b (blue line) is
shifted to a lower kinetic (high binding) energy by ~ 0.5 eV, compared to the fully screened
spectrum. This shift is constant and we used the position of the Sn 3d3/2 spectrum as a reference
to examine the effect of the local potential on the kinetic energy of the emitted electrons. The
bottom Sn 3d3/2 spectrum, shifted by 8 eV to higher kinetic energy was measured after applying
exactly – 8 eV potential to the electrode (the Sn 3d5/2 peak has moved outside the displayed
energy window). Such a shift involves the entire XPS spectrum and therefore the position of
every core level peak (i.e. Sn 3d, 4d used in this study) in the energy window of the analyzer can
be tuned electrostatically. The potential of SPEM for analyzing the charge distribution and
transport properties of working nanodevices, based on the demonstrated in Fig. 4b correlation
between the applied potential and kinetic energy of the emitted photoelectrons, is best illustrated
by the SPEM images in Fig.4 (c-e). They show percolating nanowire network measured under
three different biasing conditions. The energy window of 7.8 eV was tuned around the Sn 4d
emission of non-biased nanowires, which appear as bright stripes in Fig. 4c. On the left and right
side of the image the metal electrodes are also distinct due to the intense secondary electron
emission. The biasing of the right or left electrode (Fig.4 d, e) results in very strong attenuation
of the Sn 4d signal (due to energy shift beyond the energy window) only from certain SnO2
wires, which depends on the biased electrode. By close inspection of the images in Fig. 4c-d one
can identify some wires with rather good conductivity, which remain visible and confined points
(nods) of the high resistivity where nearly all potential drops and the entire group of nanowires
disappear from the image up to the junction point. As can be expected by biasing (Fig. 4e) of the
opposite electrode a complementing part of the network becomes visible again up to the same
In summary: using several powerful spectromicroscopy and microscopy surface science tools,
namely X-PEEM, SPEM, LEEM, LEED we obtained new information about the composition and
structure of individual metal oxide nanowires and their 2D networks. The micro-LEED patterns obtained
for the first time on individual SnO2 nanowires confirm the (010) structure of the side facets and the
preservation of the single crystal orientation of the “kinked” nanostructures. Micro- XPS and XANES
measured from the individual nanowires and nanobelts reveal surface stoichiometry, coordination and
oxidation state of the nanostructures similar to their macroscopic counterparts. Being extremely important
for electronic screening effects, apparently 100 nm size domain is still too large to induce measurable size
effects. The potential of SPEM to explore electron transport in the working chemiresistor device
composed of mats of the percolating nanowires has allowed us to identify the presence of “electro active
“elements which control the transport trough the entire percolating nanowire network. This approach
opens new avenues to explore the surface structural and electronic properties of individual nanostructures
as a function of nanostructure size, morphology, type of adsorbate, temperature. In addition, we have
shown that the method is applicable to the nanostructures wired as an active element of the functioning
nanodevice (chemiresistor in our case). The latter offers an unique experimental platform for in situ and
in real time studies in order to link fundamental surface processes taking place on the nanowire surface
with the corresponding electronic and transport properties.
Pristine single crystal quasi-1D SnO2 nanowires and nanobelts were synthesized from SnO at ca
1220 K in Ar flux (35 sccm) using the vapor-solid (VS) growth protocol. As prepared nanostructures,
with diameters and length in the range of 100-500 nm and 10-300 microns correspondingly, were
transferred from the alumina crucible and dry deposited on (i) a clean Si (111) wafer with native oxide
layer for X-PEEM/LEEM measurements and (ii) a Au foil for SPEM characterization. A model nano-
device, which allows in-situ SPEM characterization under biasing conditions, was prepared placing the
2D network of cross-linked SnO2 nanowires on a perforated Si3N4 membrane (with periodic 2 micron
holes (Aquamarijn Micro Filtration BV) and evaporating on top of the nanowire network few Au/Ti
electrodes through a shadow mask. The distance between electrodes was ca 20 mm. The average density
of the nanowires in the network was r r~0.036 mm-2, what was close to the percolation threshold r r~0.04
mm-2. This SnO2 nanowire/ Si3N4 unit was placed and wire bonded in a ceramic chip holder. Wet
processing (as those used in resist-based lithography) is avoided, because it can contaminate the surface
of the nanostructure and support and affect chemical and transport properties. The small amount of
weakly bound C-containing species on nanostructure and device surfaces are unavoidable when the
materials are in air (see survey XPS in 31 ) . They are cleaned by brief annealing in vacuum and by oxygen
plasma treatment before SPEM and X-PEEM/LEEM measurements. In particular, the sample was
annealed at 530 K for a few minutes using indirect electron bombardment heating on the back side of the
Au support. The oxygen flux was provided from a TECTRA radio frequency plasma source providing O
flux of the order of 1015 atoms cm-2 min-1.
We thank Ulrike Diebold for helpful discussions. The technical support of Clay Watts and Joshua
Cothren (SIUC) is greatly acknowledged. SIUC part of the research was supported by through ACS PRF-
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Figure 1 (a, b) Two LEEM images (FOV 20 µm) and corresponding (1x1) LEED patterns (c, d) collected from the
two areas (marked with white circles) along the individual nanostructure. The rotation of the (1x1) LEED patterns
approximately follows the bending of the nanostructure in real space. e) LEEM (FOV 6 µm, 5.5 eV) image of the
“kinked” nanostructure and (1x1) LEED patterns (45 eV) collected from different “arms”. In spite of the changed
orientation in the real space the surface structure and its orientation is preserved along the length of the
nanostructure. The inert depicts the model of the (010) facet and crystalographically equivalent growth directions.
Figure 2. XANES spectrum (green) taken at room
temperature on individual SnO2 nanowire obtained as
a ratio of the gray scale signals form XPEEM (insert)
images of the nanowire and background while
scanning the photon energy from 480 eV to 505 eV.
For comparison XAS spectra from bulk SnO (blue)
and SnO2 (red) (adapted from ) are shown. The
energy scale is offset with respect to the first
maximum of Sn 3d5/2.
Figure 3. (a) Large and small scale chemical image of individual SnO2 nanobelt obtained by monitoring the Sn
4d photoelectrons. (b) O 1s, (c) Sn 3d5/2 and (d) VB spectra taken from the SnO2 nanobelt after oxygen plasma
treatment (bottom spectrum in each panel, measured at RT) and after following annealing at 523 K (top
spectrum in each panel, measured at 523 K). The SnO band gap peak resulting from subtraction of the bottom
from the top VB spectrum is shown by the dashed line in panel (d)
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Figure 4. Spectromicroscopic access to the particularities of the electron transport and charge distribution in the
working nanowire device; a) Sketch of the experimental set-up used for SPEM characterization of the SnO2
nanowire mats under variable bias conditions. The images were measured over the energy window of Sn 4d
emission. (b) The Sn 3d3/2 and 3d5/2 spectrum taken in a selected spot of a nanowire, grounded (blue) and biased
(red) by applying potential of -8 V at the left electrode. (c)-(d) Sn 4d maps of the chemiresistor of percolating
SnO2 nanowires taken with both electrodes grounded (c), with the right (d) or left (e) electrode biased at -8 V.
The differences between the images of ‘non-biased’ and ‘biased’ structures indicate the parts of the nanowire
network that are biased, resulting in drop of the photoemission intensity due to shift of the Sn 4d emission
outside the energy window. The dependence of the electron yield on the biased electrode side is due to the
presence of highly resistive nodes (marked with white circles) which impede even potential distribution along
the nanowire network, so that only the nanostructures which remain at ground potential are bright.