JOURNAL OF APPLIED PHYSICS 108, 073708 共2010兲
Metal-insulator transition characteristics of VO2 thin films grown
on Ge„100… single crystals
Z. Yang,a兲 C. Ko, and S. Ramanathan
Harvard School of Engineering and Applied Sciences, Harvard University, Cambridge,
Massachusetts 02138, USA
共Received 25 July 2010; accepted 18 August 2010; published online 11 October 2010兲
Phase transitions exhibited by correlated oxides could be of potential relevance to the emerging field
of oxide electronics. We report on the synthesis of high-quality VO2 thin films grown on single
crystal Ge共100兲 substrates by physical vapor deposition and their metal-insulator transition 共MIT兲
properties. Thermally triggered MIT is demonstrated with nearly three orders of magnitude
resistance change across the MIT with transition temperatures of 67 ° C 共heating兲 and 61 ° C
共cooling兲. Voltage-triggered hysteretic MIT is observed at room temperature at threshold voltage of
⬃2.1 V for ⬃100 nm thickness VO2 films. Activation energies for electron transport in the
insulating and conducting states are obtained from variable temperature resistance measurements.
We further compare the properties of VO2 thin films grown under identical conditions on Si共100兲
single crystals. The VO2 thin films grown on Ge substrate show higher degree of crystallinity,
slightly reduced compressive strain, larger resistance change across MIT compared to those grown
on Si. Depth-dependent x-ray photoelectron spectroscopy measurements were performed to provide
information on compositional variation trends in the two cases. These results suggest Ge could be
a suitable substrate for further explorations of switching phenomena and devices for thin film
functional oxides. © 2010 American Institute of Physics. 关doi:10.1063/1.3492716兴
I. INTRODUCTION
Vanadium dioxide 共VO2兲 undergoes a sharp metalinsulator transition 共MIT兲 at ⬃340 K.1 Recently, it has been
observed that the MIT in VO2 can be triggered electrically2–6
in addition to the well-studied thermally induced MIT. Due
to the near-room temperature transition and voltage-triggered
MIT, a great deal of novel logic, memory, and sensor devices
are being proposed utilizing VO2.7–11 Particularly, the electrically triggered MIT in VO2 is promising for potential applications as ultrafast switches,12 Mott field effect
transistor,7,8,13 two-terminal memristive devices14 wherein
one may use the phase transition-induced resistance change
or nonlinear current-voltage 共I-V兲 characteristics as a switch.
A vertical device geometry is preferred among some of these
novel VO2 devices, particularly for using the phase transition
element to control current flow in a device or circuit.10,12 To
date, the most frequently used conducting substrate in vertical VO2 devices is heavily doped Si.4,5,15 Comparing to the
TiO2 共Refs. 16 and 17兲 and sapphire18 substrates, the crystallinity of VO2 thin films on Si substrate is compromised
due to larger lattice mismatch and formation of interfacial
oxides, however, TiO2 and sapphire are insulating and may
not be applicable as bottom contacts in VO2 based out-ofplane devices. Heavily doped Ge substrates are conducting,
and have a slightly smaller lattice mismatch with VO2 than
Si,19 but the VO2 growth on Ge substrates has not been studied so far to the best of our knowledge. Further, Ge is a high
mobility semiconductor that is attracting a great deal of interest in the transistor research community20–23 and hence is
well-suited for this study. In this paper, we report on the
a兲
Electronic mail: [email protected].
0021-8979/2010/108共7兲/073708/6/$30.00
synthesis of high-quality VO2 thin films grown on Ge substrates with the observation of both thermally triggered MIT
and room temperature voltage-triggered MIT. To the best of
our knowledge, this is the first report on growth and functional properties of VO2 films on single crystal Ge.
II. EXPERIMENTS
VO2 thin films were grown on heavily doped n+-Ge
共100兲 single-crystalline substrates 共⬃0.008 ⍀ cm resistivity兲
and reference n+-Si 共100兲 single-crystalline substrates
共0.002– 0.005 ⍀ cm resistivity兲 simultaneously with identical growth conditions using radio frequency 共rf兲 sputtering
technique with a V2O5 target. Numerous growth experiments
were carried out to identify optimal growth conditions for
good quality films. The VO2 thin film growth is quite sensitive to the oxygen gas flow rate. Figure 1 shows an example
of flow rate effect on VO2 thin films grown on sapphire
substrates using V2O5 target at growth temperature and pressure of 550 ° C and 5.0 mTorr. The total Ar and O2 gas flow
rate is 100 SCCM 共SCCM denotes cubic centimeter per
minute at STP兲. Figures 1共a兲–1共d兲 show the resistance ratio
between 25 and 100 ° C for VO2 thin films of 413 共green兲,
968 共blue兲, 1051 共red兲, and 14 共black兲 for O2 gas flow rates
of 0 SCCM, 0.5 SCCM, 0.8 SCCM, and 1.0 SCCM, respectively. The O2 gas flow rates and resistance ratios are summarized in Fig. 1共e兲. The sensitivity of VO2 properties on
synthesis conditions has been noted in earlier studies as
well24–31 and this is due to the fact that vanadium can readily
exist in multiple valence states,24–31 leading to several Magneli VnO2n−1 phases.32 Hence, great care is needed in determining optimal growth conditions that include control over
the oxygen pressure during growth to obtain VO2 films with
108, 073708-1
© 2010 American Institute of Physics
073708-2
Yang, Ko, and Ramanathan
FIG. 1. 共Color online兲 关共a兲–共d兲兴 The temperature dependence of normalized
resistance in VO2 thin films with O2 gas flow rate of 0 共green兲, 0.5 共blue兲,
0.8 共red兲, and 1.0 共black兲 SCCM using V2O5 target at 550 ° C and 5.0
mTorr. The total Ar and O2 gas flow rate is 100 SCCM. 共e兲 The plot of
resistance ratio between 25 and 100 ° C of the VO2 thin films as a function
of O2 gas flow rate.
good MIT properties. Specifically, the growth temperature,
growth pressure, gas flow rate, target gun plasma power, and
growth time were 550 ° C, 5.0 mTorr, 99.2 SCCM Ar plus
0.8 SCCM O2, 120 W, and 4 h, respectively, for the study
reported here. X-ray diffraction 共XRD兲 was performed on the
VO2 thin films in ␪-2␪ geometry using Cu K␣ of 0.1548 nm
as the x-ray source. X-ray photoelectron spectroscopy 共XPS兲
was performed using Al K␣ of 1.4866 keV as x-ray source.
Depth-dependent XPS studies were assisted with Ar ion milling. Au共200 nm兲/Ti共20 nm兲 metal contacts with a size of
100⫻ 100 ␮m2 were deposited using a Sharon e-beam
evaporator with shadow masks. The I-V characteristics were
measured by a Keithley 236 source meter integrated in a
temperature controllable probe station.
III. RESULTS AND DISCUSSION
Figure 2共a兲 shows XRD patterns of VO2 thin films
grown on Ge and Si substrates as red and green curves, respectively. The 共011兲 peak dominates in the XRD patterns of
both the VO2 thin films on Ge 共VO2-on-Ge兲 and the VO2
thin films on Si 共VO2-on-Si兲. Si and Ge possess the same
crystal structure of diamond but an ⬃4% lattice constant
mismatch.19 Hence, it is reasonable that the growth direction
of both VO2 thin films are along the same orientation 关011兴,
indicated by the XRD patterns. Several minor orientations
appear in the XRD pattern of VO2-on-Si but not VO2-on-Ge,
such as 关200兴, 关002兴, and 关210兴, indicating qualitatively that
the VO2-on-Si is less textured than VO2-on-Ge. Figure 2共b兲
shows the Gaussian fits 共blue solid lines兲 to the 共011兲 XRD
J. Appl. Phys. 108, 073708 共2010兲
FIG. 2. 共Color online兲 共a兲 XRD patterns of VO2 thin films grown on Ge
共red兲 and Si 共green兲 substrates. 共b兲 Gaussian fits 共blue solid lines兲 of 共011兲
XRD peaks of VO2-on-Ge 共red symbols兲 and VO2-on-Si 共green symbols兲.
The FWHM of the 共011兲 XRD peaks of VO2-on-Ge and VO2-on-Si is 0.50°
and 0.70°, respectively.
peaks of VO2-on-Ge 共red symbols兲 and VO2-on-Si 共green
symbols兲. The full-width-at-half-maximum 共FWHM兲 of the
共011兲 XRD peaks of VO2-on-Ge and VO2-on-Si is 0.50° and
0.70°, respectively, indicating VO2-on-Ge has slightly better
crystallinity. The average grain size of VO2-on-Ge and
VO2-on-Si is estimated to be ⬃17 and ⬃12 nm, based on
Scherrer equation 关grain size d ⬃ 共0.9␭ / FWHM cos ␪兲 with
␭ and ␪ of the x-ray wavelength and diffraction angle in the
XRD measurements33兴. The slight 共011兲 XRD peak shift in
VO2-on-Ge comparing to VO2-on-Si 共2␪ = 27.78° versus
27.95°兲, arises from the slightly smaller compressive strain
in VO2-on-Ge than that in VO2-on-Si due to the larger lattice
constant of Ge than Si.19 The VO2-on-Ge shows a slightly
less compressive strain 共by ⬃0.6%兲 than VO2-on-Si.
Figures 3共a兲 and 3共b兲 show depth-dependent XPS and
high-resolution XPS spectra of VO2-on-Ge. Figures 4共a兲 and
4共b兲 show similar XPS spectra taken from VO2-on-Si. The
curves are intentionally vertically shifted for clarity. The bottom XPS curves are taken from near the VO2 thin film surface, while the top XPS curves are near the VO2 – Ge 共or
VO2 – Si兲 interface. In the high-resolution XPS spectra, we
focus on O 1s, V 2p1/2, and V 2p3/2 peaks. The peak positions and shapes of V 2p3/2 are used to analyze the valence
states for vanadium.24–31 Initially, the dominating valence of
vanadium near the VO2 thin film surfaces on both Ge and Si
substrates is +4 共V 2p3/2 at 515.8⫾ 0.2 eV for +4
valence兲,24–31 indicating stoichiometric VO2. With the surface VO2 sputtered off by ion milling, a shoulder peak shows
up at lower energy side of the +4 V 2p3/2 peak, indicating
073708-3
Yang, Ko, and Ramanathan
FIG. 3. 共Color online兲 共a兲 Depth-dependent XPS spectra of VO2 thin film on
Ge substrate. 共b兲 High-resolution depth-dependent XPS spectra of VO2 thin
film around O 1s and V 2p regime. The curves are vertical shifted for
clarity. The bottom XPS curves are from near the VO2 thin film surface,
while the top XPS curves are near the VO2-Ge substrate interface.
the existence of a slight mixture valence of vanadium in the
film, which is very close to the +2 V valence position
共V 2p3/2 at 513.7⫾ 0.2 eV for +2 valence兲.24–31 In the
VO2-on-Ge, this shoulder peak persists until the interface
between VO2 and Ge 关Fig. 3共b兲兴, however, the shoulder peak
eventually dominates and becomes a separate peak near the
interface between VO2 and Si 关Fig. 4共b兲兴, representing a
lower valence state of vanadium. This indicates that the
VO2-on-Si shows a more evident off-stoichiometry than
VO2-on-Ge near the interface region. Figure 5共a兲 shows an
example of how the relative fraction of each valence state of
vanadium is extracted from the fitting analyses. The V 2p3/2
XPS peak 共in the high-resolution XPS spectra兲 was deconvoluted to two V 2p3/2 peaks with +4 and +2 valences. The
ratio of the areas underneath the two deconvoluted peaks
共after subtraction of the Shirley baseline background兲 gives
the relative fraction of the mixture. Figure 5共b兲 shows detailed analyses results of the relative fraction of +4 valence
vanadium in the VO2-on-Ge 共red兲 and VO2-on-Si 共green兲.
The VO2-on-Ge shows a high fraction of +4 near the surface
共⬎80%兲. Both VO2-on-Ge and VO2-on-Si show good stoichiometry 共⬃70%兲 in center region of the films 共⬃10 nm
below the surface and ⬃20 nm above the interface兲, while
the VO2-on-Ge shows a slightly higher fraction of +4 valence. When the VO2 thin films reach the near-interface regions, the fraction of +4 valence in the VO2-on-Si decreases
much faster than that of the VO2-on-Ge. This could potentially arise from interfacial reaction with the silicon substrate
owing to its propensity to form silicon oxides while germa-
J. Appl. Phys. 108, 073708 共2010兲
FIG. 4. 共Color online兲 共a兲 Depth-dependent XPS spectra of VO2 thin film on
Si substrate. 共b兲 High-resolution depth-dependent XPS spectra of VO2 thin
film focusing on O 1s and V 2p regime. The curves are intentionally vertical shifted for clarity. The bottom XPS curves are from near the VO2 thin
film surface, while the top XPS curves are near the VO2-Si substrate
interface.
nium oxide is less stable from thermodynamic
considerations.34,35 The error bars in Fig. 5共b兲 arise from the
fitting uncertainties of the inaccuracy. As noted earlier, stoichiometry control is particularly difficult for VO2 growth
due to presence of Magneli phases32 in this system and ease
of vanadium to take on several possible oxidation states.
Slight off-stoichiometry with mixture of several valence
states of vanadium is common, especially when the Si or
glass substrates are used instead of TiO2 or sapphire
substrates.24–31 Our results are consistent with these prior
reports and a slight degradation in resistance ratio across the
MIT is to be expected when compared to single-crystalline
films on lattice-matched insulating substrates or bulk single
crystals.
Figure 6共a兲 shows the temperature dependence of the
normalized in-plane resistance 关R共T兲 / R共20 ° C兲兴 of the
VO2-on-Ge substrate 共red circles兲 and VO2-on-Si 共green
circles兲. The resistance of the VO2 thin film at each temperature point is determined by the reciprocal slope of an inplane linear I-V scan between two probes. 共The linear I-V
curves indicate Ohmic contacts36 between the probes and the
VO2 thin film.兲 The solid and open circles represent temperature ramping up and down, respectively. Nearly three orders
of magnitude resistance change is observed from 20 to
100 ° C in the VO2-on-Ge, but less than two orders change in
VO2-on-Si, indicating a higher quality VO2-on-Ge than
VO2-on-Si, consistent with the XPS and XRD results. Using
Gaussian fittings on the derivative logarithmic plot of
073708-4
Yang, Ko, and Ramanathan
FIG. 5. 共Color online兲 共a兲 An example of fitting on the V 2p3/2 XPS peak
and deconvolution into two peaks representing vanadium +4 and +2 valences, showing how the relative fraction of each valence state of vanadium
is estimated. 共b兲 Fraction of the +4 valence vanadium atoms in VO2-on-Ge
共red兲 and VO2-on-Si 共green兲 along the film thickness direction.
FIG. 6. 共Color online兲 共a兲 Temperature dependence of the normalized resistance of the VO2-on-Ge 共red circles兲 and VO2-on-Si 共green circles兲. The
solid and open circles represent temperature ramping up and down, respectively. 共b兲 The MIT temperature TC of the VO2-on-Ge is determined to be
67.3 ° C and 60.8 ° C during temperature ramping up and ramping down,
respectively, using Gaussian fits on the derivative logarithmic plots. 共c兲 The
MIT temperature TC of the VO2-on-Si is determined to be 67.3 ° C and
58.6 ° C during temperature ramping up and ramping down, respectively,
using Gaussian fits on the derivative logarithmic plots.
J. Appl. Phys. 108, 073708 共2010兲
FIG. 7. 共Color online兲 共a兲 Plot of logarithmic resistance 关Ln共R兲兴 vs reciprocal temperature 共1 / T兲 of VO2-on-Ge 共red兲 and VO2-on-Si 共green兲. 共b兲 The
activation energy Ea of VO2-on-Ge 共red兲 and VO2-on-Si 共green兲 at hightemperature conducting state. 共c兲 The activation energy Ea of VO2-on-Ge
共red兲 and VO2-on-Si 共green兲 at low-temperature insulating state.
R-T,15,37,38 the MIT temperature TC of VO2-on-Ge is determined to be 67.3 ° C and 60.8 ° C during temperature ramping up and ramping down, respectively, as shown in Fig.
6共b兲, while 67.3 ° C and 58.6 ° C for VO2-on-Si during temperature ramping up and ramping down, respectively, as
shown in Fig. 6共c兲. The VO2-on-Ge shows 2.2 ° C narrower
MIT width 共6.5 ° C versus 8.7 ° C兲 than the VO2 thin film on
Si. The minor difference in the transition temperature and
MIT width between VO2-on-Ge and VO2-on-Si is possibly
due to variation in the faction of +4 valence 关other valence
vanadium oxide phases can distort the MIT 共Ref. 31兲兴 and
strain39 within the films as discussed earlier. However, the
observed thermal-MIT hysteresis window width and transition temperatures in both VO2 thin films show consistent
trends with previous reported values.4,5,15,38
Figure 7共a兲 shows the plot of logarithmic resistance
关Ln共R兲兴 versus reciprocal temperature 共1 / T兲, which is employed for the activation energy 共Ea兲 analyses. The activation
energy is estimated by the slope of a linear fitting of Ln共R兲
over 共1 / kBT兲, as Ln共R兲 = b + Ea / kBT, where R, T, kB, and b
are resistance, temperature 共in unit K兲, Boltzmann constant,
and intercept. Figures 7共b兲 and 7共c兲 show the activation energy of VO2-on-Ge 共red兲 and VO2-on-Si 共green兲 in lowtemperature “insulating” and high-temperature “metallic”
states. The activation energy of VO2-on-Ge is 193⫾ 7 meV
and 149⫾ 4 meV for low- and high-temperature states, respectively. The reason that the high-temperature activation
energy is not significantly smaller than that of the lowtemperature is likely due to the mixed valence states of the
VO2 thin films. Once the temperature ramps up across the
MIT transition temperature, only the VO2 phase in the film
undergoes transition into the metallic state. In other words,
even at high temperatures the film is not entirely metallic;
073708-5
J. Appl. Phys. 108, 073708 共2010兲
Yang, Ko, and Ramanathan
hence the activation energy is not significantly smaller than
its low-temperature insulating state. In a previous study,4 for
example, a VO2 thin film with nearly four orders of magnitude resistance change in thermal-MIT, which is closer to the
ideal VO2, shows a smaller high-temperature activation energy than the low-temperature one 共80–90 meV versus 230–
250 meV兲. For the VO2-on-Si sample, again due to mixture
of valences and smaller order of resistance change across
MIT than VO2-on-Ge, the high-temperature activation energy is not distinct from the low-temperature one, considering the error bars. This likely indicates that the hightemperature state is a mixture of metallic VO2 regions and
some insulating phases 共from other valence vanadium oxide
phases兲. Similar results on uniform temperature dependence
of resistance on both sides of the MIT were also observed
previously in off-stoichiometric VO2 films grown by
electron-beam evaporation.40 We do note here that the estimated activation energies are comparable to previous
studies.41–43 Kristensen41 prepared VO2 using powder sintering method and found an activation energy of 240 meV from
temperature dependence of the electrical conductivity analyses. Sahana et al.42 grew VO2 thin films on glass substrates
by chemical vapor deposition and the temperature-dependent
resistance measurements indicated that the activation energy
of the VO2 thin films ranged from 180–260 meV depending
on growth conditions.
Figure 8共a兲 shows the I-V characteristics of VO2-on-Ge
measured in vertical geometry between 0 and 5 V at 25 ° C
共red兲, 60 ° C 共green兲, and 90 ° C 共blue兲. Voltage-triggered
MIT is observed at 25 ° C with threshold voltage at ⬃2.1 V.
The inset in Fig. 8共a兲 shows the magnified region of the
resistive hysteresis window of the I-V curves at 25 ° C, with
scan directions marked with arrows. The width of the voltage
triggered MIT window is ⬃1 V. Considering the ⬃100 nm
thickness of the VO2 thin film, the threshold electric-field for
triggering MIT is approximately 2.1⫻ 107 V / m at 25 ° C,
also in good agreement with studies on both VO2 planar
junctions on sapphire and capacitor-type devices on
silicon.2–7 In the hysteresis curves, some “staircase” discontinuity shapes are observed, which is likely due to grain
boundary or domain related contributions to MIT.44 Similar
phenomena were also observed in nanoscale VO2 planar
junctions by Sharoni et al.,45 who concluded that the
thermal-MIT in VO2 occurs through a series of staircase
shape avalanches induced by impact ionization due to grain
boundary effects. The effects were seen primarily in nanoscale junctions and hence reasonable to expect in our measurements since we are studying voltage-triggered phase
transition across an ⬃100 nm film. Each staircase step may
correspond to a grain/domain boundary associated change
共slight differences in the insulating state conductivity lead to
shifts in the critical voltage needed to initiate the transition
and hence a step-like shape emerges兲. Electrically triggering
MIT studies by conducting atomic force microscopy but in
regions smaller than a single grain size, shows free of staircase steps in the I-V curves,46 which supports the analyses
here and those in Ref. 45.
Figure 8共b兲 shows the Poole–Frenkel 共PF兲 plot 关Ln共I / V兲
versus Sqrt 共V兲兴 of the I-V curves shown in Fig. 8共a兲. At
FIG. 8. 共Color online兲 共a兲 I-V characteristics of the VO2-on-Ge measured in
vertical geometry between 0 to 5 V at 25 ° C 共red兲, 60 ° C 共green兲, and
90 ° C 共blue兲. The electric-field-triggered MIT is observed at 25 ° C. The
inset shows the magnified region of the resistive hysteresis window of the
I-V curves at 25 ° C, with scanning directions marked with arrows. 共b兲 PF
plot of the I-V curves 关Ln共I / V兲 vs Sqrt 共V兲兴. Three regions, metallic, Ohmic,
and PF dominating are defined. The two blue bold lines in the PF region are
for visual guidance, indicating the dominating of PF conduction mechanism
in that region.
large voltages 共⬎ ⬃ 3.1 V兲, the VO2 thin film shows metallic
behavior over a broad range of temperature. At low voltages
共⬍ ⬃ 0.7 V兲, the VO2 thin film only shows metallic behavior
at high temperatures, but behaves as insulator at lower temperatures. At intermediate bias voltages and below thermalMIT temperature, the PF mechanism dominates the conduction, indicated by the linear dependence of Ln共I / V兲 over Sqrt
共V兲. The two bold blue lines in the PF region is used for
visual guidance for the linear dependences. Three regions,
namely metallic, “Ohmic” and “PF dominating” are denoted
in Fig. 8共b兲, showing the dominant conduction mechanism in
each region. At either high-temperature or large applied voltages, the VO2 enters metallic state due to the thermally or
electrically triggered MIT. At low temperatures and small
bias voltage, Ohmic conduction behavior is predominant in
VO2. At low-temperature and intermediate bias voltage, the
PF mechanism dominates. Nonstoichiometric defects, such
as oxygen vacancies in VO2, may serve as localized traps
required for the PF emission. Prior to onset of metallic state,
the localized states assist the PF conduction under electricfield bias.4,5,47
IV. SUMMARY
VO2 thin films were grown on Ge共100兲 single crystal
substrates using rf sputtering. Thermally triggered MIT with
nearly three orders of magnitude resistance change with an
073708-6
Yang, Ko, and Ramanathan
MIT width of ⬃6 ° C is observed in VO2 thin films grown on
Ge. Hysteretic voltage-triggered MIT is observed at room
temperature at a critical voltage of ⬃2.1 V with a hysteresis
window of ⬃1 V in VO2 thin films grown on Ge. The structural and electrical measurements suggest that Ge may be a
suitable substrate for further explorations of phase transitionbased oxide electronics utilizing MITs.
ACKNOWLEDGMENTS
This work was supported by ONR Grant N00014-10-10131 and AFOSR Grant FA9550-08-1-0203.
F. J. Morin, Phys. Rev. Lett. 3, 34 共1959兲.
G. Stefanovich, A. Pergament, and D. Stefanovich, J. Phys.: Condens.
Matter 12, 8837 共2000兲.
3
B. G. Chae, H. T. Kim, D. H. Youn, and K. Y. Kang, Physica B 369, 76
共2005兲.
4
C. Ko and S. Ramanathan, Appl. Phys. Lett. 93, 252101 共2008兲.
5
D. Ruzmetov, G. Gopalakrishnan, J. Deng, V. Narayanamurti, and S. Ramanathan, J. Appl. Phys. 106, 083702 共2009兲.
6
D. Ruzmetov, G. Gopalakrishnan, C. Ko, V. Narayanamurti, and S. Ramanathan, J. Appl. Phys. 107, 114516 共2010兲.
7
H. T. Kim, B. G. Chae, D. H. Youn, S. L. Maeng, G. Kim, K. Y. Kang, and
Y. S. Lim, New J. Phys. 6, 52 共2004兲.
8
S. Hormoz and S. Ramanathan, Solid-State Electron. 54, 654 共2010兲.
9
M. J. Lee, Y. Park, D. Suh, E. H. Lee, S. Seo, D. C. Kim, B. Jung, B. S.
Kang, S. E. Ahn, C. B. Lee, D. H. Seo, Y. K. Cha, I. K. Yoo, J. S. Kim,
and B. H. Park, Adv. Mater. 共Weinheim, Ger.兲 19, 3919 共2007兲.
10
T. H. Yang, R. Aggarwal, A. Gupta, H. H. Zhou, R. J. Narayan, and J.
Narayan, J. Appl. Phys. 107, 053514 共2010兲.
11
A. Gupta, R. Aggarwal, P. Gupta, T. Dutta, R. J. Narayan, and J. Narayan,
Appl. Phys. Lett. 95, 111915 共2009兲.
12
A. Cavalleri, C. Tóth, C. W. Siders, J. A. Squier, F. Ráksi, P. Forget, and
J. C. Kieffer, Phys. Rev. Lett. 87, 237401 共2001兲.
13
C. Zhou, D. M. Newns, J. A. Misewich, and P. C. Pattnaik, Appl. Phys.
Lett. 70, 598 共1997兲.
14
T. Driscoll, H. T. Kim, B. G. Chae, M. Di Ventra, and D. N. Basov, Appl.
Phys. Lett. 95, 043503 共2009兲.
15
C. Ko and S. Ramanathan, J. Appl. Phys. 106, 034101 共2009兲.
16
Y. Muraoka and Z. Hiroi, Appl. Phys. Lett. 80, 583 共2002兲.
17
J. Lu, K. G. West, and S. A. Wolf, Appl. Phys. Lett. 93, 262107 共2008兲.
18
K. Okimura, J. Sakai, and S. Ramanathan, J. Appl. Phys. 107, 063503
共2010兲.
19
In monoclinic VO2, lattice constant a = 0.5753 nm. The lattice constants
of Si and Ge are 0.5431 nm and 0.5646 nm, respectively.
20
S. Verdonckt-Vandebroek, E. F. Crabbe, B. S. Meyerson, D. L. Harame, P.
1
2
J. Appl. Phys. 108, 073708 共2010兲
J. Restle, J. M. C. Stork, and J. B. Johnson, IEEE Trans. Electron Devices
41, 90 共1994兲.
21
M. Ieong, B. Doris, J. Kedzierski, K. Rim, and M. Yang, Science 306,
2057 共2004兲.
22
M. L. Lee, E. A. Fitzgerald, M. T. Bulsara, M. T. Currie, and A.
Lochtefeld, J. Appl. Phys. 97, 011101 共2005兲.
23
T. Maeda, Y. Morita, and S. Takagi, Appl. Phys. Express 3, 061301
共2010兲.
24
G. A. Sawatzky and D. Post, Phys. Rev. B 20, 1546 共1979兲.
25
J. Mendialdua, R. Casanova, and Y. Barbaux, J. Electron Spectrosc. Relat.
Phenom. 71, 249 共1995兲.
26
M. Ghanashyam Krishna, Y. Debauge, and A. K. Bhattacharya, Thin Solid
Films 312, 116 共1998兲.
27
M. Demeter, M. Neumann, and W. Reichelt, Surf. Sci. 454–456, 41
共2000兲.
28
G. Silversmit, D. Depla, H. Poelman, G. B. Marin, and R. De Gryse, J.
Electron Spectrosc. Relat. Phenom. 135, 167 共2004兲.
29
N. Alov, D. Kutsko, I. Spirovová, and Z. Bastl, Surf. Sci. 600, 1628
共2006兲.
30
A. Romanyuk and P. Oelhafen, Thin Solid Films 515, 6544 共2007兲.
31
D. H. Youn, H. T. Kim, B. G. Chae, Y. J. Hwang, J. W. Lee, S. L. Maeng,
and K. Y. Kang, J. Vac. Sci. Technol. A 22, 719 共2004兲.
32
J. B. Goodenough, J. Solid State Chem. 3, 490 共1971兲.
33
A. Patterson, Phys. Rev. 56, 978 共1939兲.
34
W. Kaiser, P. H. Keck, and C. F. Lange, Phys. Rev. 101, 1264 共1956兲.
35
L. M. Terman, Solid-State Electron. 5, 285 共1962兲.
36
L. J. Mandalapu, Z. Yang, and J. L. Liu, Appl. Phys. Lett. 90, 252103
共2007兲.
37
D. Brassard, S. Fourmaux, M. Jean-Jacques, J. C. Kieffer, and M. A. E.
Khakani, Appl. Phys. Lett. 87, 051910 共2005兲.
38
D. Ruzmetov, K. T. Zawilski, V. Narayanamurti, and S. Ramanathan, J.
Appl. Phys. 102, 113715 共2007兲.
39
T. Kikuzuki and M. Lippmaa, Appl. Phys. Lett. 96, 132107 共2010兲; B.
Lazarovits, K. Kim, K. Haule, and G. Kotliar, Phys. Rev. B 81, 115117
共2010兲.
40
R. G. Mani and S. Ramanathan, Appl. Phys. Lett. 91, 062104 共2007兲.
41
I. K. Kristensen, Mater. Res. Bull. 9, 1677 共1974兲.
42
M. B. Sahana, M. S. Dharmaprakash, and S. A. Shivashankar, J. Mater.
Chem. 12, 333 共2002兲.
43
J. Li and N. Yuan, Appl. Surf. Sci. 233, 252 共2004兲.
44
A. Frenzel, M. M. Qazilbash, M. Brehm, B. G. Chae, B. J. Kim, H. T.
Kim, A. V. Balatsky, F. Keilmann, and D. N. Basov, Phys. Rev. B 80,
115115 共2009兲.
45
A. Sharoni, J. G. Ramírez, and I. K. Schuller, Phys. Rev. Lett. 101,
026404 共2008兲.
46
J. Kim, C. Ko, A. Frenzel, S. Ramanathan, and J. E. Hoffman, Appl. Phys.
Lett. 96, 213106 共2010兲.
47
P. P. Boriskov, A. A. Velichko, A. L. Pergament, G. B. Stefanovich, and D.
G. Stefanovich, Tech. Phys. Lett. 28, 406 共2002兲.