Letterpubs.acs.org/NanoLett

True Vapor−Liquid−Solid Process Suppresses Unintentional Carrier
Doping of Single Crystalline Metal Oxide Nanowires
Hiroshi Anzai,† Masaru Suzuki,† Kazuki Nagashima,† Masaki Kanai,† Zetao Zhu,† Yong He,†
Mickael̈ Boudot,† Guozhu Zhang,† Tsunaki Takahashi,† Katsuichi Kanemoto,‡ Takehito Seki,§Naoya Shibata,§ and Takeshi Yanagida*,††Institute for Materials Chemistry and Engineering, Kyushu University, 6-1 Kasuga-Koen, Kasuga, Fukuoka 816-8580, Japan
Department of Physics, Osaka City University, 3-3-138 Sugimoto, Sumiyoshi-ku, Osaka 558-8585, Japan§
Institute of Engineering Innovation, The University of Tokyo, 2-11-16 Yayoi, Bunkyo-ku, Tokyo 113-8656, Japan‡S Supporting Information*

ABSTRACT: Single crystalline nanowires composed ofsemiconducting metal oxides formed via a vapor−liquid−
solid (VLS) process exhibit an electrical conductivity even
without an intentional carrier doping, although thesestoichiometric metal oxides are ideally insulators. Suppressing
this unintentional doping eﬀect has been a challenging issue
not only for metal oxide nanowires but also for various
nanostructured metal oxides toward their semiconductor
applications. Here we demonstrate that a pure VLS crystal
growth, which occurs only at liquid−solid (LS) interface,substantially suppresses an unintentional doping of single
crystalline SnO2 nanowires. By strictly tailoring the crystal growth interface of VLS process, we found the gigantic diﬀerence ofelectrical conduction (up to 7 orders of magnitude) between nanowires formed only at LS interface and those formed at both LS
and vapor−solid (VS) interfaces. On the basis of investigations with spatially resolved single nanowire electrical measurements,
plane-view electron energy-loss spectroscopy, and molecular dynamics simulations, we reveal the gigantic suppression of
unintentional carrier doping only for the crystal grown at LS interface due to the higher annealing eﬀect at LS interface comparedwith that grown at VS interface. These implications will be a foundation to design the semiconducting properties of various
nanostructured metal oxides.
KEYWORDS: Metal oxide nanowires, vapor−liquid−solid growth, unintentional carrier doping, crystal growth interfaceS

nonstoichiometry exhibits the exotic physical properties ofmetal oxides, including catalytic, memristive properties, andothers.16−27 On the contrary, the unintentional doping eﬀect
has not been a serious issue for Si, Ge, and III−V nanowires,
indicating the inherent diﬀerence between ionic metal oxide
nanowires and covalent semiconductor nanowires on theoccurrence of unintentional doping during crystal growth.
Thus, controlling the electrical properties and the unintentional
doping eﬀect is an inherent and challenging issue forsemiconducting metal oxide nanowires.
Here, we demonstrate that a pure VLS crystal growth inwhich a solid phase is only formed at LS interface, substantiallysuppresses an unintentional doping of single crystalline SnO2
nanowires, which had been impossible by any existing metaloxide nanowire growth methodologies. To achieve the strict
growth control, we utilized “material ﬂux window principle”,
which describes a relationship between a material ﬂux and aingle crystalline nanowires composed of semiconducting
metal oxides, including SnO2, ZnO, In2O3 and others, haveshown their great promises for applications in a variety of ﬁelds
ranging from energy conversion and electronics to biotechnology due to their unique physical properties and robustness inwater and air that are hardly attainable to conventional
semiconductors.1−6 However, tailoring the electrical conductivity of these metal oxide nanowires has been very diﬃcult due
to the occurrence of unintentional doping during crystal growth
process.7−9 In fact, most nanostructured semiconducting metaloxides in general are electrically conductive even without any
intentional doping,10,11 although these stoichiometric oxides
should be insulators.12 Suppressing such unintentional doping
eﬀect has been a challenging and long-standing issue not only
for above single crystalline semiconducting metal oxide
nanowires but also for various nanostructured metal oxides
toward their semiconductor applications, which require thecontrol of carrier concentration and the bipolarity.13−15
Nonstoichiometric properties (e.g., a crystal imperfection) ofmetal oxides due to the strong ionicity are closely related to theoccurrence of such unintentional doping.7−11 Note that such
© 2017 American Chemical Society

Received: March 31, 2017
Revised: June 14, 2017
Published: July 3, 2017
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Figure 1. (a) SEM image of tapered SnO2 nanowires. (b) SEM image of untapered SnO2 nanowires. (c) Single nanowire electrical measurement
device for tapered SnO2 nanowires. (d) Single nanowire electrical measurement device for untapered SnO2 nanowires. (e) Comparison between
tapered- and untapered nanowires on the electrical measurement (current−voltage data). (f) Statistical data of electrical resistivity when varying Sn
ﬂux and nanowire diameter.

nucleation probability of solid at a LS or a VS interface.28 On
the basis of this knowledge, we can control the contribution of
two crystal growth interfaces (LS and VS) for single crystalline
SnO2 nanowire formation by varying the vapor ﬂux density.
Interestingly, the resistance diﬀerence between nanowires
formed only via LS interface and nanowires formed via both
LS and VS interfaces is found to be gigantic up to 7 orders ofmagnitude (ranged from kΩ to TΩ), although these two
nanowires are grown under the same conditions except for thevapor ﬂux density. Spatially resolved single nanowire electricalmeasurements, plane-view scanning transmission electron
microscopy (STEM), electron energy-loss spectroscopy(EELS) analysis, and molecular dynamics (MD) simulations
reveal that the unintentional doping via a crystal imperfection
can be drastically suppressed at LS interface.
This work is triggered by the discovery of the gigantic
diﬀerence between tapered and untapered SnO2 nanowires on
the electrical conductivity, as shown in Figure 1. Figure 1a,bshows the scanning electron microscopy (SEM) images offabricated SnO2 nanowires when we vary the Sn metal ﬂux(ranged from 10.4 to 27.2 × 1017 cm−2 s−1) during nanowire
growth. Except for the Sn ﬂux density, these SnO2 nanowires
were grown under the same conditions, including growthtemperature (750 °C), oxygen partial pressure (10−3Pa), andAu catalyst thickness (∼1 nm). The details of nanowire growth
can be seen in Methods. Nanowires grown under relatively
higher Sn ﬂux (27.2 × 1017 cm−2 s−1) showed the tapered
nanowire shape, whereas such tapering was not observable fornanowires grown under relatively lower Sn ﬂux (10.4 × 1017
cm−2 s−1). Figure 1c,d shows the SEM images of the four-probe

devices to measure the electrical properties of a single nanowire
for both tapered and untapered SnO2 nanowires. Figure 1eshows the comparison between these two nanowires on theelectrical properties (current−voltage data). The resistance of
untapered SnO2 nanowire was over 1 TΩ, which is surprisingly
7 orders of magnitude higher than that of tapered nanowire(93.0 kΩ). To further reveal the statistical diﬀerence between
tapered and untapered SnO2 nanowires on the electricalconduction, Figure 1f shows the distribution of measured
electrical resistivity when varying the Sn ﬂux and the nanowire
diameter. As clearly seen in the ﬁgure, there is a signiﬁcantdiﬀerence between tapered and untapered SnO2 nanowires on
the single nanowire electrical resistivity. The electrical resistivity
showed the diﬀerence of 5 orders of magnitude in the nanowire
diameter range around 50 nm when varying the Sn ﬂux.
Considering the fact that these tapered and untapered
nanowires were grown under thermodynamically similar
conditions except for the Sn ﬂux, this observed diﬀerence isgigantic. Thus, varying slightly Sn ﬂux strongly aﬀects not only
the tapering of nanowire shape but also the electrical properties
of SnO2 nanowires during VLS growth.
Above tapering phenomena of nanowires grown via VLS
process can be interpreted in terms of the competition between
a crystal growth at VS interface on the nanowire sidewall (VSgrowth) and a crystal growth at LS interface (LS growth).29
This is also consistent with so-called “material ﬂux window”,
which describes the concept of VLS nanowire growth within alimited material ﬂux range, where a nucleation preferentiallyoccurs only at the LS interface.28 In this principle, increasing
the Sn ﬂux should tend to promote the VS growth, resulting in
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Figure 2. (a) Sn ﬂux dependence on SnO2 nanowire morphology (diameter) and SnO2 ﬁlm growth rate. (b) Growth time dependence of nanowire
diameter. These SnO2 nanowires were grown under same the conditions, including growth temperature (750 °C), oxygen partial pressure (10−3 Pa),
and Au catalysts thickness (∼1 nm).

Figure 3. (a) Schematic illustration of multielectrodes devices to measure the spatial electrical distribution of a tapered SnO2 nanowire. (b) SEM
image of multielectrodes devices to measure the spatial electrical distribution of a tapered SnO2 nanowire. (c) Spatial distribution of 4-probe
electrical resistance for a tapered SnO2 nanowire with calculated values for spatially homogeneous model. (d) Spatial distribution of temperaturedependent 4-probe electrical resistance for a tapered SnO2 nanowire.

the tapering of nanowire shape, which is consistent with above
our results. To ensure this tapering eﬀect in our system, we
measured the morphology of nanowires when varying the Sn
metal ﬂux. Figure 2a shows the Sn ﬂux dependences on both
the diameter value near the tip and the maximum diametervalue at the bottom. Above 15 × 1017 cm−2 s−1 of the ﬂux, thetapering occurred, where the maximum diameter value tends toincrease whereas the diameter near the tip was independent of
the ﬂux variation. To more clearly reveal the role of VS growthon these diameter data, we measured the Sn ﬂux dependence
on the VS ﬁlm growth rate in the absence of Au catalysts, as
shown in Figure 2a. These ﬁlms were grown under the same
growth conditions of nanowires. Below 15 × 1017 cm−2 s−1 ofthe ﬂux, there was almost no VS ﬁlm growth, and the ﬁlm
growth rate increased as the ﬂux increased above 15 × 1017
cm−2 s−1. Clearly, the Sn ﬂux dependence of VS growth canwell explain the tapering phenomena of the present SnO2

nanowires in terms of the competition between VS growth and
LS growth. To further reveal the absence of VS growth for
untapered SnO2 nanowires, we measure the growth timedependence of diameters for untapered nanowires because thediameter of untapered nanowires should increase with thegrowth time if the untapered SnO2 nanowires are formed atboth LS and VS interface. The results are shown in Figure 2b.
As can be seen, the diameter for untapered nanowires is almost
consistently independent of the growth time. Thus, the crystal
growth at VS interface for untapered nanowires is almost
negligible, which is also consistent with the results of Figure 2a.
All our experimental results consistently support the absence of
crystal growth at VS interface for untapered SnO2 nanowires.
These results highlight that we can strictly control the crystal
growth interface including VS interface and/or LS interface
during the VLS nanowire growth of SnO2 by varying the Sn
metal ﬂux. Thus, overall experimental trends of Figure 1 and
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Figure 4. (a) Plane-view STEM image of tapered SnO2 nanowire. (b) Plane-view STEM image of untapered SnO2 nanowire. (c) Spatial distributionof EELS spectra for both tapered SnO2 nanowire and untapered SnO2 nanowire with reference EELS spectra for SnO, SnO2, and SnO2−δ.

relatively short spatial distance within 3 μm for this single
nanowire transport measurement, the observed 3 orders
diﬀerence of resistance values and the diﬀerent temperaturedependence for diﬀerent electrode gaps were rather gigantic. In
the ﬁrst and second scenarios, the diameter dependence of
nanowire resistance should be approximately proportional to 1/
diameter2. This is because these two scenarios assume theconductive inner layer, whose cross-sectional area increases
with increasing the nanowire diameter. The third scenario’s
diameter dependence of nanowire resistance can be stronger,
which depends on the resistivity diﬀerence between the crystal
grown at VS interface (VS crystal) and the crystal grown at LSinterface (LS crystal), and also the cross-sectional area ratio of
each crystal. To examine this, we measured the spatial diametervariation within a single nanowire in SEM image. Figure 3c also
shows the comparison between experimental data and expected
diameter dependence of homogeneous model. The experimental diameter dependence on the nanowire resistance did
not conform to 1/diameter2 of spatially homogeneous model,
and the diameter dependence is much stronger than 1/
diameter2. In addition, all spatially resolved single nanowire
measurements showed that the nanowire resistivity tended to
be lower as the measured spatial position tended to be awayfrom the tip, where VS crystal tended to be dominant. Thisimplies the signiﬁcant role of outer VS crystal on the electricaltransport properties. Note that the similar gigantic resistancevariation within a single tapered SnO2 nanowire was also
observed even when we varied the oxygen partial pressures topromote the VS growth instead of the variation of Sn ﬂux, as
seen in Supporting Information S2. Thus, these spatiallyresolved resistance measurement data within the single tapered
SnO2 nanowire consistently support the third scenario, whichassumes the conductive outer VS crystal and the rather
insulative inner LS crystal, to explain the gigantic diﬀerence of
nanowire resistance when varying the Sn metal ﬂux in Figure
1e.
Next, we question what is essentially diﬀerent between LS
crystal and VS crystal. Because X-ray diﬀraction and ultraviolet−visible (UV−vis) light absorption spectra data can notreveal the spatial diﬀerence between the two crystals (seeFigure 2 imply that enhancing the VS crystal growth at
nanowire sidewall (i.e., the tapering) seems to strongly impact
the electrical transport properties of a single SnO2 nanowire.
Here we consider the following three possible scenarios toexplain the gigantic diﬀerence between the tapered anduntapered SnO2 nanowires on the electrical transport properties in Figure 1e,f. The ﬁrst scenario is based on the nanowire
diameter diﬀerence, which explains the nanowire resistancevariation in terms of the cross-sectional area diﬀerence byassuming the spatially homogeneous electrical conductivity.
Second scenario assumes the spatially inhomogeneous
conduction model based on the presence of insulating
depletion layer near the nanowire surface, which explains theresistance increase for the smaller nanowires. Third possible
scenario assumes the conductive shell layer formed via VSinterface and the insulative core crystals formed via LSinterface. The illustration of these three models is shown in
Supporting Information S1. To examine the ﬁrst scenario, wemeasure the electrical resistivity of a single nanowire whenvarying Sn ﬂux, as shown in Figure 1f. As can be seen in Figure
1f, increasing the Sn ﬂux value consistently decreases theresistivity value. Thus, the observed Sn ﬂux dependence on theelectrical transport properties cannot be interpreted solely interms of the diameter variation via tapering. This excludes theﬁrst scenario based on only the nanowire diameter diﬀerencewith the spatially homogeneous electrical conductivity.
To further validate above three scenarios, we measured thespatial distribution on the electrical transport properties oftapered SnO2 nanowires by using multielectrodes, as shown in
Figure 3a,b. The tapering eﬀect (i.e., VS growth eﬀect) on theelectrical properties must be less signiﬁcant as the measured
spatial position approaches to the tip of nanowire. Figure 3cshows the spatial distribution data of 4-probe resistances within
the single tapered SnO2 nanowire. Figure 3d shows thetemperature-dependent electrical conductivity. As can be seen,
there was the strong spatial dependence of nanowire resistance,
which was ranged from 104 to 107 Ω. As the measured spatial
position tends to be away from the tip of tapered nanowire, theresistance tended to decrease and the temperature dependence
tended to be metallic from semiconducting. Considering the4701

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Figure 5. (a) Snap shots of crystal growth MD simulations for both LS interface and VS interface in the presence of impurity atoms. (b) Timeevolution data of crystallized atom numbers for both LS interface and VS interface. (c) Escaping probability data of impurity atoms from a solid
phase into a liquid or vapor phase during crystal growth as a function of bonding number. (d) Schematic illustration of an inherent diﬀerence
between LS interface and VS interface on an unintentional carrier doping during crystal growth.

unintentional doping of SnO2 during crystal growth, although
the exact origin of unintentional doping in SnO2 has been along-standing issue and controversial.31,32 As to the relationshipbetween such crystal imperfections-oxygen vacancies and theelectrical conductivity for SnO2−x, we measure the electricalresistivity of SnO2−x ﬁlms when varying the oxygen partialpressure during ﬁlm formation; see the details in Supporting
Information S7. Results imply that there is a certain correlation
between crystal imperfections−oxygen vacancies and theelectrical conductivity of SnO2−x. These STEM-EELS results
highlight that the tapered SnO2 nanowire is composed of theouter VS crystal layer, which contains crystal imperfections,
including oxygen deﬁciencies, and the inner more stoichiometric LS crystal core.
Finally, we question why there is a signiﬁcant diﬀerence
between LS crystal and VS crystal on the electrical conductivity
via the crystal imperfection and/or the impurity incorporation
dynamics during a VLS nanowire growth of SnO2. Because it israther diﬃcult to experimentally observe the impurity
incorporation dynamics, we qualitatively examine this issue by
performing molecular dynamics (MD) simulation. The details
of MD simulations can be seen in Methods and Supporting
Information S6. In these MD simulations, we intentionallyemploy impurity atoms to examine the impurity incorporation

Supporting Information S3 and S4), we performed plane-view
STEM-EELS analysis. The details of STEM-EELS measurements can be seen in Methods. Figure 4a,b shows the planeview annular dark ﬁeld (ADF) STEM images of tapered anduntapered SnO2 nanowires. Both plane-view STEM images
consistently showed the single crystallinity, and the diﬀerent
cross-sectional nanowire shapes are due to the occurrence ofVS growth on the sidewall of nanowires. To observe thediﬀerence between LS crystal and VS crystal, we have
performed spatially resolved EELS analysis, and the spatial
spectra data of both tapered and untapered SnO2 nanowires areshown in Figure 4c. For comparison, the reference spectra30 of
SnO2, SnO2−δ, and SnO are shown. For untapered nanowire,
there was no spatial dependence of the O−K edge spectra data,
which well agrees with the reference spectra data ofstoichiometric SnO2. On the other hand, for tapered nanowire
the O−K edge EELS spectra data exhibited the spatial
dependence. The spectra data near the outer layer tends to
be closer to the reference data of oxygen-deﬁcient SnO2−δ,whereas the spectra data of inner core is very consistent withthe reference data of stoichiometric SnO2 (see more details in
Supporting Information S5). This crystal imperfection,
including oxygen deﬁciencies and related other dopant
incorporations, seems to play an important role on the4702

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tures (e.g., bipolarity and pn junctions) of various functional
semiconducting metal oxide nanostructures.
We have shown that a pure VLS crystal growth only at LSinterface substantially suppresses an unintentional doping of
single crystalline SnO2 nanowires. By strictly tailoring the
crystal growth interface of VLS process, we found the gigantic
electrical conduction diﬀerence (up to 7 orders of magnitude)
between nanowires formed only at LS interface and thoseformed at both LS and VS interfaces. Spatially resolved single
nanowire electrical measurements revealed the occurrence ofpreferential unintentional doping only for the crystal growth at
VS interface. Plane-view STEM-EELS analysis clariﬁed the
crystal imperfection of outer-layer of nanowires, formed at VSinterface, which was not observable for nanowires formed at LSinterface. MD simulations qualitatively explain why such
unintentional doping preferentially occurs only at VS interface
and not at LS interface in terms of the annealing eﬀect
enhanced at LS interface. These implications will be afoundation to further tailor the semiconducting properties
not only of single crystalline metal oxide nanowires but also of
various functional metal oxide nanostructures.
Methods. Fabrication of SnO2 Nanowires. SnO2 nanowires were grown onto a sapphire Al2O3(110) substrate(Electronics and Materials) by pulsed laser deposition method,
as reported elsewhere.36−38 Prior to nanowire growth, Au layer(∼1 nm) was deposited onto the substrate by utilizing direct
current sputtering (SANYU Electron, SC-701HMC). A laser
ablation with ArF excimer laser (λ = 193 nm, Coherent,
COMPex Pro) was performed to supply the vapor source.
SnO2 tablet (99.99% pure, Kojundo Chemical Laboratory) wasused as the target. For the nanowire growth, the chamber wasﬁrst evacuated to be less than 5 × 10−5 Pa, then the oxygen/argon mixture gas (99.9999% pure) was introduced bycontrolling the oxygen partial pressure in the range of 10−4−
10 Pa. The total pressure of 10 Pa was kept in order to preventthe catalyst consumption during the nanowire growth.39 Before
the laser ablation, the substrate temperature was heated up to
750 °C within 20 min and kept constant during the nanowire
growth. The repetition rate, the laser energy, and the target-tosubstrate distance were 10 Hz, 40 mJ, and 30 mm, respectively.
The ﬂux of vapor species was controlled by varying the energy
density (mJ/mm2) of laser ablation. The duration of nanowire
growths was 30−300 min. After the nanowire growth, thesubstrate temperature was cooled down to room temperature
within 30 min. The Sn ﬂux was quantitatively evaluated via ﬁlm
deposition at room temperature in 10 Pa O2 atmosphere. Aftermeasuring the thickness of the obtained thin ﬁlm using asurface proﬁler (KLA Tencor, Alpha-Step IQ), the material ﬂux
was estimated using the density of the element.
Nanowire Characterizations. The morphology of obtained
SnO2 nanowires was evaluated by a ﬁeld-emission scanningelectron microscopy (FESEM, JEOL, JSM-7610F) at accelerating voltage of 15−30 kV. The diameter and the length ofnanowires were statistically evaluated from 100 nanowires. We
employed a scanning transmission electron microscopy(STEM, JEOL, JEM-ARM300F) equipped with electron energy
loss spectroscopy (EELS, GIF Quantum), at accelerating
voltage of 300 kV. A Cu microgrid covered with lacey C ﬁlm
was used to support the nanowire samples for the STEM
measurements. The probe size was approximately 0.06 nm.
Acquisition time of EELS spectra was 45 s for each pixel.
During the acquisitions, the probe scanned inside pixel. In
order to correct energy drift due to instability of the

dynamics into LS crystal or VS crystal. In the following
simulations, the relative interaction parameter of impurity ε′ isset to be ε′ = 0.8 (<1), where the incorporation of impurity
atoms into the solid phase destabilizes the solid structure.
Figure 5a shows the snapshots of performed MD simulations
for both LS crystal and VS crystal. As illustrated in Figure 5a,
the impurity incorporation into the solid phase tends to be
more diﬃcult for LS interface growth when compared with VS
interface growth. To understand this diﬀerence on the impurity
incorporation dynamics in terms of the growth dynamics, we
calculated the time tracking data of the number of atoms
participating the crystal growth in MD simulations, as shown in
Figure 5b. The time tracking data of both LS crystal and VS
crystal were shown for the comparison. As the crystal growthcontinues, the number of crystallized atoms increases for both
crystals. One of the clear diﬀerence between LS crystal and VS
crystal on the time tracking data is a ﬂuctuation in the numberof crystallized atoms. The number of crystallized atoms at LSinterface is more ﬂuctuated than that at VS interface,highlighting that the growing layer at the solid surface ismore frequently reconstructed at LS interface. Figure 5c shows
the diﬀerence between LS interface and VS interface on theescaping probability of impurity atoms from the solid phase
into vapor or liquid phase during crystal growth. At LS
interface, an impurity atom once incorporated into the solid
phase can be more eﬀectively ejected into liquid phase,compared with the ejection of impurity atoms to vapor phase at
VS interface. In consequence, excessive impurity atoms are wellreleased from the solid phase, and more pure solid structure canbe grown at LS interface, which is in contrast to the VS
interface where unintentional crystal imperfections remain.
This eﬀect is very similar to an annealing eﬀect when we anneal
the samples to improve the crystallinity and purity.33−35 In
general, atoms are more frequently exchanged between liquidand solid phases, when compared to that between vapor andsolid phases. This is because the eﬀective energy per atom islower in the liquid than in the vapor phase when the liquid has
high aﬃnity with the atoms. This kind of annealing eﬀect
enhanced by liquid phase and the improvement of thecrystallinity should generally occur at LS interface. Thus,
these MD simulations qualitatively demonstrate that anunintentional doping of VLS nanowires can be suppressed byintentionally forming the crystal only at LS interface without
the crystal at VS interface, as illustrated in Figure 5d.
The present results highlight that we can tailor the previously
unavoidable unintentional carrier doping of nanostructured
semiconducting metal oxides by utilizing the pure VLS crystal
growth only at LS interface. The resistivity values of present
well-deﬁned SnO2 nanowires are almost 3−4 orders ofmagnitude higher than those of previously reported single
crystalline SnO2 nanowires (see Supporting Information S8).
For this well-deﬁned VLS process, the metal catalysts play arole of ﬁltering to exclude impurity atoms from fabrication
environments. In other words, it is possible to fabricate highpurity semiconducting metal oxide nanostructures even underrelatively “dirty” and cheap fabrication circumstances compared
to conventional semiconductor fabrication technologies, which
require ultraclean and expensive growth processes. MD
simulations predict that this strategy should be applicable forvarious semiconducting metal oxide nanostructures if thecrystal growth is designed to occur only at LS interface via
VLS process. The present knowledge can be extended tofurther tailor the semiconducting properties and heterostruc4703

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microscope, the zero-loss peak was simultaneously obtained ateach core-loss acquisition. Because the original reference
spectra in Figure 4 are not calibrated, they were oﬀset toalign the oxygen ﬁrst peak. The smoothed spectra wereobtained by weighted average method with 1 eV window. X-ray
diﬀraction (XRD, Philips, X’Pert) was employed to evaluate theaveraged crystallinity of nanowires on the substrate. Ultraviolet−visible (UV−vis) absorption spectroscopy (Shimadzu,
UV-2400) was performed in the wavelength range of 190−900
nm. The absorption spectra of the SnO2 nanowires wereobtained by subtracting the background signals from thesapphire substrate.
Electrical Transport Measurements. The electrical transportmeasurement of SnO2 nanowires was conducted with a form of
single nanowire device, which was constructed onto SiO2/Si
substrate by 30 kV electron beam (EB) lithography
technique.40,41 For the SiO2/Si substrate, thermally oxidized
100 nm thick SiO2 was capped on the degenerated p-type Si(100) substrate (resistivity ∼1−5 × 10−3 Ωcm). EB lithography
process was performed by utilizing ZEP520A-7 (ZEON) as a
resist. First, the nanowires grown onto sapphire substrate weredispersed into isopropanol by sonication with frequency of 28
kHz for 1 min. Then 1 μL of the nanowire suspension was
dropped onto the SiO2/Si substrate with prepatterned
microelectrode. EB lithography was performed to bridge
between the nanowires and the microelectrodes. For theelectrical contact, Pt was deposited via a custom-ordered RF
magnetron sputtering (SEINAN) at 50 W in Ar 0.3 Pa. The Pt
thickness was controlled to be about 150 nm. Finally, the resist
capped with Pt ﬁlm was lifted-oﬀ by N,N-dimethylformamide
followed by the cleaning in acetone. The electrical transportmeasurements were performed in the probe station (Lake
Shore Cryotronics, Model TTPX) connected with a semiconductor parameter analyzer (Keithley, 4200SCS) at temperature range of 100−300 K in vacuum condition (<10−2 Pa).
Molecular Dynamics (MD) Simulations. The molecular
dynamics (MD) simulation can express time-evolution of
motion (position and translation speed) for each of particles
which interact with each other through short-range LennardJones (L-J) interaction.42−44 In this simulation system, fourtypes of particles (A, B, A′, and C) were utilized to model
vapor−liquid−solid (VLS) growth process. Here, A and C canmake AC compound in solid phase that expresses an oxide
formation, A′ is an impurity atom, and B is assigned as solvent
liquid atom. L-J type potential is used as two-body interaction
potential except that the interaction between A−A and C−C.
More details of calculations and employed formulas can be seenin Supporting Information S7.

■

Takeshi Yanagida: 0000-0003-4837-5701
Notes

The authors declare no competing ﬁnancial interest.

■

ACKNOWLEDGMENTS
This work was supported by CREST of Japan Science andTechnology Corporation (JST). M.B was supported by
ImPACT. K.N. and H.Y. were supported by KAKENHI Grants(26706005, 16H00969, 15K13288, 15H03528, and 26220908).
N.S. acknowledges a support from the Grant-in-Aid for
Scientiﬁc Research on Innovative Areas “Nano Informatics”(Grant 25106003) from JSPS. A part of this work was
conducted at the Research Hub for Advanced Nano Characterization, the University of Tokyo, under the support of“Nanotechnology Platform” (Project No.12024046) by the
MEXT. This work was performed under the Cooperative
Research Program of “Network Joint Research Center for
Materials and Devices” and the MEXT Project of “Integrated
Research Consortium on Chemical Sciences”.

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ASSOCIATED CONTENT

S Supporting Information*

The Supporting Information is available free of charge on the
ACS Publications website at DOI: 10.1021/acs.nanolett.7b01362.
Additional information and ﬁgures (PDF)■

REFERENCES

AUTHOR INFORMATION

Corresponding Author*E-mail: yanagida@cm.kyushu-u.ac.jp.
ORCID

Katsuichi Kanemoto: 0000-0001-7437-3940
Naoya Shibata: 0000-0003-3548-5952
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DOI: 10.1021/acs.nanolett.7b01362
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