Download www.rsc.org/nanoscale Volume 5 | Number 16 | 21 August 2013

www.rsc.org/nanoscale
Volume 5 | Number 16 | 21 August 2013 | Pages 7077–7640
ISSN 2040-3364
PAPER
Arbiol, Xiong et al.
Solution phase van der Waals epitaxy
of ZnO wire arrays
Nanoscale
View Article Online
PAPER
Cite this: Nanoscale, 2013, 5, 7242
View Journal | View Issue
Solution phase van der Waals epitaxy of ZnO wire
arrays†
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Yue Zhu,a Yong Zhou,a Muhammad Iqbal Bakti Utama,a Marı́a de la Mata,b
Yanyuan Zhao,a Qing Zhang,a Bo Peng,a Cesar Magen,c Jordi Arbiol*bd
and Qihua Xiong*ae
As an incommensurate epitaxy, van der Waals epitaxy allows defect-free crystals to grow on substrates
even with a large lattice mismatch. Furthermore, van der Waals epitaxy is proposed as a universal
platform where heteroepitaxy can be achieved irrespective of the nature of the overlayer material and
the method of crystallization. Here we demonstrate van der Waals epitaxy in solution phase synthesis
for seedless and catalyst-free growth of ZnO wire arrays on phlogopite mica at low temperature. A
Received 22nd April 2013
Accepted 8th May 2013
unique incommensurate interface is observed even with the incomplete initial wetting of ZnO onto the
substrate. Interestingly, the imperfect contacting layer does not affect the crystalline and optical
properties of other parts of the wires. In addition, we present patterned growth of a well-ordered array
DOI: 10.1039/c3nr01984e
with hexagonal facets and in-plane alignment. We expect our seedless and catalyst-free solution phase
www.rsc.org/nanoscale
van der Waals epitaxy synthesis to be widely applicable in other materials and structures.
Introduction
The ever-growing research interest in heteroepitaxy, or the
growth of one crystalline material on another, is motivated by
substrate engineering, heterojunction devices, and device
integration.1 Typically, utilization of heteroepitaxy is focused on
lattice-matched or nearly lattice-matched systems. Meanwhile,
lattice-mismatched heteroepitaxial thin lm growth that is
coherent and dislocation-free could be achieved only below a
certain critical thickness, which is usually on the order of a few
nanometers.2 In the case of nanowires as well, there exist critical
diameter and height factors,3 but the requirement for defectfree growth is less stringent compared to thin lms due to
a
Division of Physics and Applied Physics, School of Physical and Mathematical
Sciences, Nanyang Technological University, Singapore 637371. E-mail: Qihua@ntu.
edu.sg
b
Institut de Ciència de Materials de Barcelona, ICMAB-CSIC, Campus de la UAB,
08193 Bellaterra, Catalonia, Spain. E-mail: [email protected]
c
Laboratorio de Microscopı́as Avanzadas (LMA), Instituto de Nanociencia de Aragon
(INA) - ARAID and Departamento de Fisica de la Materia Condensada, Universidad
de Zaragoza, 50018 Zaragoza, Spain
d
Institució Catalana de Recerca i Estudis Avançats (ICREA), 08010 Barcelona,
Catalonia, Spain
e
Division of Microelectronics, School of Electrical and Electronic Engineering, Nanyang
Technological University, Singapore 639798
† Electronic supplementary information (ESI) available: Wire array diameter
distribution curve (Fig. S1); supplementary power spectra (Fig. S2); SEM images
of a ZnO lm grown on GaN (Fig. S3); a ZnO nanowire array on ZnO seeded Si
(Fig. S4); a pattern formed by laser interference lithography with the
corresponding growth of an ordered array (Fig. S5); a SEM image of ZnO wires
hydrothermally grown on Bi2Se3 akes (Fig. S6). See DOI: 10.1039/c3nr01984e
7242 | Nanoscale, 2013, 5, 7242–7249
lateral relaxation.4 Such a limitation in lattice-mismatched
systems originates from the prerequisite that epitaxial crystalline overlayer (epilayer) materials form covalent bonding with
the substrate. However, it is also possible that only a weak van
der Waals interaction exists at the epilayer–substrate interface,
if the surface of the substrate is a van der Waals surface, i.e.,
when the surface lacks dangling bonds and is thus chemically
inert.5 Utilizing this principle, growth of dislocation-free crystals over a lattice-mismatched substrate could be realized with
the mechanism called van der Waals epitaxy (vdWE). It is
important to point out that although the interface is mediated
by a weak van der Waals force, vdWE is still able to reach the
equilibrium state in terms of surface energy. The most direct
evidence is the observation of preferred crystalline orientations
of the epilayer with respect to the substrate, justifying the
classication of vdWE as an epitaxy mechanism.6
The application of vdWE was pioneered by Koma and coworkers in 1980 for planar lm growth by molecular beam
epitaxy.7–9 The capability of vdWE to produce non-planar
nanostructures was reported only very recently, such as in II–VI
semiconductors and alloyed nanowire arrays and nanotripods
on muscovite mica by a vapor transport process,10–13 indium
arsenide nanowires on graphene by metal–organic chemical
vapor deposition,14 and topological insulator bismuth selenide
nanoplates on uorophlogopite mica by vapor phase transport.15 It should be noted that all of the above examples can be
classied as vapor phase vdWE at a high growth temperature,
whereas no attention is paid to the case of solution phase
synthesis of non-planar structures although van der Waals
epitaxy has been indeed invoked in the solution synthesis of a
This journal is ª The Royal Society of Chemistry 2013
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Paper
Langmuir–Blodgett lm.16 In solution phase epitaxy, the epilayer material is synthesized from chemical reaction in solution, in which soluble precursors precipitate to nucleate and
grow on the substrate. Solution phase epitaxy is also distinct
from liquid phase epitaxy, a well-established physical method
in which a very high-purity supersaturated melt is used, from
which the crystalline semiconductor material is precipitated
without involving chemical reaction.17 However, solution phase
epitaxy has not been well established experimentally, especially
for the case of heteroepitaxy. The major obstacle is that energy
provided during solution phase synthesis is usually not high
enough to overcome the high interfacial energy to achieve heteroepitaxy. Therefore, the reported solution phase epitaxy
encompasses only limited examples in which a small lattice
constant mismatch largely reduces the interfacial energy to
allow conventional covalent heteroepitaxy, such as the ZnO–
GaN system (1.9% mismatch)18 and TiO2–uorine-doped SnO2
system (2% mismatch).19
In this paper, we report the realization of vdWE of non-planar
structures from solution phase synthesis, utilizing hydrothermal synthesis of ZnO wires as a model system and phlogopite mica as a substrate. Phlogopite mica (KMg3AlSi3O10(OH)2)
possesses a layered structure and can be easily cleaved perfectly
in one direction producing thin sheets or akes, which are
transparent to translucent and are also exible and elastic.
Being chemically inert, its surface satises the requirement of
van der Waals epitaxy.5 Surprisingly, van der Waals epitaxy can
be achieved in low-temperature solution synthesis. We reveal
the Volmer–Weber growth mechanism of the array in solution
phase synthesis, deviating slightly from that in the vapor phase.
As one example of the advantage of the current solution phase
synthesis, we show the controlled growth of an ordered array on
a patterned mica substrate by laser interference lithography,
which fosters future device fabrication potential. In the end, we
present the excellent optical properties of the as-synthesized
arrays, despite the low reaction temperature.
Results and discussions
A large area and high density vertically aligned ZnO wire array
can be obtained on phlogopite mica, as revealed from scanning
electron microscopy (SEM) observations (Fig. 1a and b) on the
sample synthesized in a glass reagent bottle. The top view SEM
image (Fig. 1c) shows an in-plane alignment of the hexagonal
faceting among the individual ZnO wires, similar to our
previous report for vapor phase synthesis.11 This observation of
in-plane alignment strongly indicates that the wires are grown
epitaxially, which is further conrmed to be vdWE later by the
epitaxy analysis as shown in Fig. 2. In addition, many of the
hexagonal facets are of similar size, indicating a relatively
narrow diameter distribution of the array (Fig. S1 in the ESI†)
although no seed/catalyst or patterning has been employed.
There are indeed wires with larger diameters, and their top
facets deviate from ideal hexagonal shape. This observation will
be discussed later in the part addressing the growth mechanism. The crystal orientation of the as-grown array is shown by
X-ray diffraction (XRD) measurements (Fig. 1d). Besides the
This journal is ª The Royal Society of Chemistry 2013
Nanoscale
(00L) basal planes of phlogopite mica (monoclinic, space group:
C2/c), only the (0002) peak from wurtzite ZnO (space group:
P63mc) can be observed. The absence of other peaks from ZnO
indicates that the array is oriented well along the h0001i c-axis
direction.
Although the buffer layer is commonly observed in conventional epitaxial growth of ZnO nanowires,20,21 it is completely
absent in the vapor phase vdWE.11,13 In contrast, the solution
phase synthesis reported herein produces a discontinuous layer
that is observable beneath the ZnO body of the wire (Fig. 2a). A
survey of the ZnO–mica interface was performed by scanning
transmission electron microscopy (STEM) in high angle annular
dark eld (HAADF) imaging mode (Fig. 2b) to examine such a
distinct structure at the base of the ZnO wire. As the HAADF
offers Z-contrast, the lighter phlogopite mica will appear darker
than the heavier (higher averaged atomic number) ZnO, allowing a rapid identication of the interface and the contacting
region. We observed that the wetting of ZnO onto the mica
substrate is not complete: there is alternation between contacting regions (indicated by arrows) and void-like regions at the
heterointerface. It should be noted that the thickness of such a
non-uniform layer is around 20 nm for the wire that we
observed. Above this initial thin layer, the ZnO region shows a
uniform scanning prole corresponding to the perfect hexagonal body.
We proceed to directly visualize the contacting interface of
ZnO–mica with high resolution transmission electron microscopy (HRTEM). The resulting image (Fig. 2c) shows that the ZnO
wire is monocrystalline without any crystalline defects or
dislocations even from the rst few atomic layers. The buffer
layer, usually highly defective, is absent at the ZnO–mica
interface (as marked with a dark blue line). The zoom-in
HRTEM images (Fig. 2d) show the perfect lattice fringes of ZnO
and phlogopite mica, each at a distance of about 10 nm from
the interface. The perfection of the ZnO crystal validates that
solution phase growth is effective in producing the well-crystallized non-planar structure. However, the ZnO region right
above the interface appears uneven (as marked with a yellow
dashed line in Fig. 2c), having brighter contrast than the region
above it. We believe that such an uneven interface corresponds
to the void-like regions in Fig. 2b. The ZnO crystal near the
interface exhibits contrast modulation that could be attributed
to the presence of strain. Although there is no obvious dislocation, the lattices are somewhat distorted. We ascribe such
distortions partially to the strain created during focused ion
beam (FIB) preparation for cross-sectional HRTEM imaging.
The presence of the Pt and Au protective layers deposited
around the ZnO for FIB preparation may create Moiré fringes
when overlapping with ZnO (see the Experimental section for
details on FIB preparation). More importantly, these distortions
could also originate from the coalescence effect during the
growing process, which will be discussed later in the growth
mechanism (Fig. 3).
The indexed power spectra (Fig. 2e) from the fast Fourier
transform (FFT) of the HRTEM images show the ZnO, the
interface and the mica regions as recorded separately. The
phase and lattice spacing in the c-axis from the power spectra of
Nanoscale, 2013, 5, 7242–7249 | 7243
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Nanoscale
Paper
Fig. 1 Morphological characterization of the as-synthesized vertically aligned ZnO wire array on the phlogopite mica substrate. (a) Pseudocolored SEM image with 45
tilted view in low magnification. Scale bar: 0.1 mm. The region colored in blue contains high density of wires, demonstrating the capability of synthesis over a wide area.
The region in red is phlogopite mica the top surface of which is exfoliated, showing the layered structure of the substrate. (b) SEM image with 45 tilted view in high
magnification, showcasing the high density of wires. Scale bar: 10 mm. (c) Top view SEM image. Scale bar: 10 mm. The arrow is a guide to show the orientation and inplane alignment of the hexagonal faceting of the wires. (d) XRD pattern collected in the q–q configuration, revealing the orientation of the as-grown array. The peaks
are indexed in blue for ZnO and in red for phlogopite mica.
ZnO (Fig. 2e, upper) and mica (Fig. 2e, lower) are consistent with
those measured by XRD (Fig. 1d). From the color-combined
power spectrum of the ZnO–mica interface (Fig. 2e, middle), the
epitaxial relationship of the ZnO–mica can be determined as
0] ZnO||(001)[010]mica. We notice that the arrange(0001)[112
ment and spacing of spots in the power spectra of ZnO at the
heterointerface are nearly identical to those away from the
heterointerface. The lattice mismatch, which is measured to be
6.94% in the present zone axis, is still observable even aer
the heteroepitaxial growth (i.e., the epitaxy is incommensurate).
Despite the relatively small lattice mismatch, the Bragg reec00) ZnO planes at the heterointertion spot of the in-plane (11
face is not forced to coincide with that of the (200)mica which is
also in the in-plane direction (Fig. 2e, middle). Similarly, the
00] ZnO||[100] mica
observation of the heterointerface in the [11
zone axis also reveals a large lattice mismatch of around 64.9%
which is not accommodated by the in-plane ZnO lattices (Fig. S2
in the ESI†). In other words, the ZnO lattices are already relaxed
even from the heterointerface, where the lattices exhibit a bulk
value of lattice parameters with negligible mismatch-related
strain. The relaxation of ZnO lattice could also be inferred from
the isolated Bragg diffraction spots instead of streaks. We thus
conclude that the ZnO–mica interface complies with the characteristics of incommensurate vdWE growth. Such a conclusion
is reasonable because we are using a layered phlogopite mica as
the substrate: the surface of layered materials is typically free of
7244 | Nanoscale, 2013, 5, 7242–7249
dangling bonds and is thus unlikely to allow covalent bonding
required in the conventional heteroepitaxial growth.7 The
generalizability of vdWE in both vapor and solution phases is
therefore established.
The non-uniform contacting layer at the ZnO–mica interface
is of particular interest as it is related to the very initial growth
stage of the wires. The voids observed at the heterointerface
indicate the existence of regions where the nucleation species
are not sufficiently adsorbed to cause partial wetting.22 In
comparison, vdWE in the case of vapor phase synthesis of ZnO
nanowire arrays on muscovite mica achieved complete wetting
throughout the base of individual nanowires.11 In the current
solution synthesis, wetting behavior shis to partial wetting,
which could be attributed to an elevated interfacial energy
resulted from precursor molecules and ions, their decomposition products and even water molecules present in the system.23
In addition, the low reaction temperature used in the solution
phase does not favor the surface migration, which could ll any
gaps among different nucleation islands. Up to this point, we
can picture the rst stage of array growth in solution phase:
nucleation species are rst adsorbed onto the mica surface to
form nucleation islands. Due to the synergic effect of all factors
in the solution environment, the wetting of each individual
island is incomplete, resulting in small contact areas of 5 to
10 nm. Above the contacting layer, however, the nucleation
produces a perfect single crystalline structure (Fig. 2c).
This journal is ª The Royal Society of Chemistry 2013
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Paper
Nanoscale
Fig. 2 The characteristics of van der Waals epitaxy in ZnO wire-phlogopite mica from solution phase synthesis. (a) SEM image of the base of a wire, showing the
existence of a discontinuous layer. Scale bar: 500 nm. (b) HAADF-STEM image of the ZnO–mica heterointerface, revealing the nonuniformity of the contacting layer with
voids among the contact regions (indicated by arrows). Scale bar: 50 nm. (c) Cross-sectional HRTEM image of the region surrounding the ZnO–mica heterointerface. The
interface is marked with a dark blue line. Scale bar: 10 nm. (d) Zoom-in HRTEM images of the regions as indicated by the corresponding colored squares and arrows in
(c). Scale bar: 2 nm. (e) Power spectra (from FFT of the HRTEM images) of ZnO (pseudocolored in red), mica substrate (in green) and the interface (color-combined) of the
regions shown in (d). The perfect relaxation of the ZnO lattices at the heterointerface is consistent with the features of the incommensurate vdWE mechanism.
On the substrate that was taken out from the reaction solution shortly heated in a convection oven, only tiny wires were
observed (Fig. 3a). The diameter of those initial stage wires is
much smaller than that of the nal array. Therefore, we
conclude that the initial growth pattern in solution phase vdWE
follows the Volmer–Weber (VW) model, where three-dimensional islands grow on the substrate and coalesce into a single
structure at a later stage.24 The SEM images of samples taken
Fig. 3 Structural characterization and investigation of the growth mechanism. SEM images of samples collected from the reaction solution at (a) 1 h, (b) 4 h and (c) 12 h
(45 tilted view, scale bar: 1 mm). (d) HAADF-STEM images of a single wire with two separated legs. Left: image of the whole wire. Middle: zoom-in image of the base
region of the wire, exhibiting the separation. Scale bar: 50 nm. Right: high resolution imaging of one of the legs in the region inside the red square (scale bar: 2 nm) and the
power spectrum of the same region. (e) Left: ADF STEM image of a single wire from the merging of two wires. Scale bar: 0.2 mm. Right: surface plot of the pink square in the
left image of (e) showing the coalescence of two wires, each having a hexagonal cross-section. (f and g) HRTEM images of the bottom region of ZnO wires, showing (f) a
merged region with an apparent gap and (g) a region with the gap closed. Scale bars: 5 nm. In all images, white arrows indicate the growth direction of the wire.
This journal is ª The Royal Society of Chemistry 2013
Nanoscale, 2013, 5, 7242–7249 | 7245
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Nanoscale
out aer a longer reaction time (Fig. 3b and c) suggest that wires
also grow laterally (i.e., increase in cross-sectional area), despite
the growth rate being much slower than the axial growth (i.e.,
increase in height). Such a fast growth in the c-axis of ZnO is
commonly observed and is attributed to the high reactivity of
the polar +(0001) surface.25 Moreover, we made several interesting observations in the SEM images of temporal evolution of
the ZnO array. At the very initial stage (Fig. 3a), two or more tiny
wires were observed close to each other, occupying a base area
that is approximately equal to that of a nal ZnO wire. In the
next growth stage (Fig. 3b), clusters with two or more wires were
also observed besides individual wires. As the growth
approaches the nal stage (Fig. 3c), most of the wires were
single-standing with larger diameter, though some wires with
much smaller diameter were also present. We also observed
wires that appear to have two heads (e.g., at the bottom right of
Fig. 3c), possibly the result of reseparation between two wires
that were previously merged. These observations further
conrm the VW growth model where large wires are produced
through the coalescence of multiple nucleation events. Thus we
can picture the second stage of array growth: tiny wires grow
from nucleation islands in the rst stage due to the growth
preference along the c-axis (Fig. 3a). Two or more adjacent wires
may coalesce together due to lateral growth, forming a complete
wire. However, the gaps among these wires at the bottom (the
wetting region) persist during the growth process, leaving the
discontinuous layer (Fig. 2a) aer the reaction. One would also
expect the formation of even larger wires through the merging
of two or more individual distant nucleation events, which can
be considered as secondary coalescence. Indeed, these wires
correspond to the large ones shown previously (Fig. 1c), usually
possessing irregular top facets. Although these wires fail to form
a complete hexagon due to the lack of enough surrounding
nucleation sites, they still maintain the in-plane alignment due
to epitaxy.
To further support our hypothesis, we characterized the
bottom region of wires removed from the mica substrate by
ultrasonication. The HAADF-STEM images of two wires with
apparent bottom features are shown (Fig. 3d and e): one with a
large open gap and one with a closed gap, both showing the case
of two individual nucleation events coalescing together. Notice
that for the rst case, the individual wire is single crystalline as
conrmed by performing the power spectrum on the aberration
corrected HAADF STEM image (FFT, Fig. 3d inset). For the
second case, the surface plot shows that the cross-section forms
two adjacent hexagons (Fig. 3e, inset). The cross-section is
consistent with the remnant of two free-standing wires before
the coalescence. Epitaxial growth ensured that the coalescence
results in a well-crystallized manner: the wires exhibit in-plane
alignment (Fig. 1c) such that coalescing wires should have
00i planes that
identical planes to be in contact, i.e., h11
construct the side facets. The HRTEM image of the merging
point between a gap region and a coalesced region (Fig. 3f)
shows the matching of lattice fringes which belong to two
different crystals. We even observed similar matching in a
region without an obvious gap (Fig. 3g), showing the merging of
two regions by different contrast perpendicular to the growth
7246 | Nanoscale, 2013, 5, 7242–7249
Paper
direction. This merging mechanism could also give a hint on
the lattice distortions mentioned previously in Fig. 2c. If we
consider the presence of multiple nucleation events (i.e.,
multiple “legs”), when the resulting legs are in contact with
each other and merge into a single wire body, there is a low
probability for the atoms to attach perfectly. In this case, the
material would suffer certain strain until the different grain
boundaries would merge into a perfect crystal. This point thus
provides an explanation for the high strain and distortion
observed during the rst few nanometers of growth, with all the
legs competing for axial growth against each other.
In heteroepitaxy, the coalescence is possible through inplane alignment, and its extent is mainly determined by interfacial energy. In conventional covalent heteroepitaxy with small
lattice mismatch, if the substrate is single crystalline, a
continuous lm may more likely be obtained instead of singlestanding nanowires (low interfacial energy). In contrast, if the
substrate is not single crystalline or if there is no long range
ordering of crystalline domains, single standing nanowires
without in-plane alignment will be the major product. We use
single crystal GaN substrate (prepared by metal–organic
chemical vapor deposition on sapphire) and ZnO-seeded Si
(prepared by radio frequency magnetron sputtering) substrates
for comparison with the mica substrate. Under the same reaction conditions as those for ZnO wire growth on mica, a lm of
ZnO is obtained on GaN (Fig. S3a†) while a nanowire array is
obtained on Si (Fig. S4a†). We noticed that the lm of GaN is
formed through coalescence of hexagonal ZnO units as seen
from an imperfect area (Fig. S3b†). This phenomenon is similar
to that observed on the mica substrate, except the scale, which
is much larger due to the strong covalent bonding between the
GaN substrate and the ZnO overlayer. The nanowire array
formed on Si shows less ordering in terms of vertical alignment
and in-plane alignment as compared to the array on the mica
substrate. More importantly, the average diameter of the
nanowires on Si (around 100 nm) is much smaller than that of
the wires on the mica substrate. These two differences may be
explained by the random orientation of the ZnO seeds on Si:
though oriented in the c-axis direction during sputtering,26
there is no in-plane alignment of these seeds, eliminating the
possibility of merging of any adjacent nanowires. Therefore,
the as-grown nanowires show a diameter more or less close to
the grain size of sputtered seeds (Fig. S4b†).
Position-specic synthesis of a wire array is an important
step towards device integration, thus it is crucial to achieve
patterning on the mica substrate. Here, we adopt the laser
interference lithography27 to fabricate a periodic circular hole of
similar size to the diameter of previous wires. Aer the formation of patterns on the mica substrate (Fig. S5a†), the synthesis
was conducted without further modications.
As can be seen from the SEM images (Fig. 4a and b), an
ordered array was obtained. Besides the apparent ordering
following the pre-fabricated pattern, these wires exhibit more
uniform hexagonal morphology than the random array
(Fig. 1b). Similarly, the top view (Fig. 4b) reveals the ordered
arrangement of these wires on the mica substrate with
pronounced in-plane alignment observed previously for the
This journal is ª The Royal Society of Chemistry 2013
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Paper
Nanoscale
Fig. 4 Controlled growth of a well-ordered wire array on patterned mica. (a)
SEM image with a 60 tilted view in high magnification to display the wellordered wires. Scale bar: 5 mm. (b) Top view SEM image confirms the ordering and
displays the in-plane alignment of the wires. Scale bar: 5 mm.
random array (Fig. 1c). A unique feature of patterned growth is
the absence of irregularly shaped wires formed through coalescence. Due to the spatial connement of photoresist, adjacent wires were prevented from merging together. As a
consequence, the majority of the wires produced from
patterned growth exhibit a nearly ideal hexagonal shape of
similar size. Indeed, observation on the side of the ordered
array indicates that each wire is grown from an individual hole
without interference from other wires (Fig. S5b†).
We characterized the emission property of the as-synthesized array on non-patterned mica to show that these wires
retain good crystallinity despite the presence of non-uniform
contacting layer. The typical room-temperature steady-state
photoluminescence (PL) spectra of the array (Fig. 5a) were
obtained under the excitation of a solid state 355 nm laser. Two
peaks can be readily resolved, which is in good agreement with
previous studies.28,29 The narrow peak located at 387 nm can be
assigned to the near band-edge emission. Generally, the broad
band centered around 580 nm is attributed to the recombination of electrons in the oxygen vacancies with excited holes in
the valence band.30 With the increase of excitation power, the
relative intensity ratio of the broad band to the narrow peak is
decreased, supporting that the broad band emission is due to
some defect states. The red shi of the near band-edge emission
is typically attributed to the heating effect of laser illumination.31 Even under a very low excitation power condition, the
deep level emission is much weaker than the near band-edge
This journal is ª The Royal Society of Chemistry 2013
Fig. 5 Photoluminescence spectra of the as-grown ZnO array on non-patterned
mica. (a) Power dependent room temperature photoluminescence spectra. Inset:
linear UV-emission peak intensity with pumping power. (b) Temperature dependent photoluminescence spectra.
emission, indicating that our ZnO wire arrays have a low defect
density. The temperature dependent PL spectra (Fig. 5b) were
obtained at the xed pumping power of 850 mW. The spectra
maintain the shape at room temperature: the near band-edge
emission is always signicantly strong at different temperatures, with only blue-shi of the peak according to Varshni's
law.32 This thus further conrms the excellent optical properties
of the as-synthesized array.
Conclusion
In conclusion, we demonstrate the vdWE in solution phase
synthesis of the ZnO array on phlogopite mica. Different from
that in the vapor phase, the growth in this case shows a nonuniform wetting of the substrate, resulting in individual nucleation events following the Volmer–Weber mechanism. However,
perfectly shaped crystals can still be formed as wires grown from
Nanoscale, 2013, 5, 7242–7249 | 7247
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Nanoscale
individual nucleation sites can merge into a single wire. The
perfection of the as-grown arrays can be further demonstrated
by the excellent photoluminescence with relatively low defect
emission intensity. In patterned growth, an ordered array with
more uniform morphology can be obtained, with the elimination of coalescence. Besides the mica substrate, ZnO wires could
also be grown on other van der Waals surfaces, for example the
Bi2Se3 akes prepared by the scotch-tape method (Fig. S6 in
the ESI†). Our solution phase vdWE is solid evidence to prove
the universality of vdWE as an incommensurate epitaxy strategy.
In particular, we provide a potential platform on which the
convenience of solution phase synthesis could be exploited;
besides the demonstrated patterned growth herein, previous
strategies of morphology control33,34 and doping35,36 could be
directly adopted. Our work also largely extends solution phase
epitaxy that was previously limited to several systems based on
conventional heteroepitaxy. We further anticipate that such
solution phase van der Waals epitaxy may also be used for both
planar and non-planar crystal growth in a variety of materials.
Moreover, the extension into other substrates with van der
Waals surfaces and substrates with intentional dangling bonds
passivation37,38 is also promising. The excellent emission property could be harnessed to fabricate optoelectronic devices such
as LED39 and electrically pumped laser.40
Paper
Characterization
Morphology and crystallinity of ZnO arrays on phlogopite mica
were characterized by eld-emission scanning electron
microscopy (JEOL JSM-7001F) and X-ray diffraction (Bruker D8
advanced diffractometer, Cu Ka radiation). The cross-sectional
sample was prepared by a 5–30 keV Ga+ ion polishing in a
focused ion beam system (Helios 600 Nanolab), preceded by a
deposition of Au and Pt thin layers to protect the sample during
the milling process. The ZnO–mica interface was observed by
high resolution transmission electron microscopy (JEOL 2010F,
with an acceleration voltage of 200 kV). Normal scanning
transmission electron microscopy and high angle annular dark
eld scanning transmission electron microscopy were performed on an aberration corrected probe FEI Titan 60–300 keV
operated at 300 kV. The as-grown ZnO wire array sample was
also dispersed in isopropanol by ultrasonication, followed by
the drop-casting of the suspension onto a lacey carbon grid for
further observation. Optical characterization was conducted
using a spectrometer (Horiba-JY T64000) in a backscattering
conguration. The back-scattered signal was collected through
a 100 objective, dispersed by a 1800 g mm1 grating under a
single mode with a spectral resolution of 1 cm1, and recorded by a liquid nitrogen-cooled charge coupled device detector.
The laser was focused by a UV objective (40, NA: 0.4) and then
excited on samples.
Experimental section
Material synthesis
Acknowledgements
The ZnO arrays were hydrothermally synthesized via the
convenient chemical bath deposition in aqueous solutions
containing 20 mM zinc nitrate and hexamethylenetetramine.41
Generally, the average diameter of a wire array is positively
related to the precursor concentration up to a point where too
concentrated solution (above 80 mM) would produce a lm-like
structure. For a typical batch, the mixture was rigorously stirred
in a glass reagent bottle before reaction. A freshly cleaved thin
layer of phlogopite mica (V-1 Grade, SPI) was made to oat on
the reaction solution with the fresh surface facing down. The
backside of the mica substrate was sealed with poly(methyl
methacrylate) to prevent the substrate from sinking. Then the
bottle was tightly capped and was heated in a water bath at
80 C. Aer the desired reaction time was reached, the mica
substrate was taken out from the reactor, rinsed with deionized
water and blow-dried under nitrogen ow.
Q. X. gratefully acknowledges the strong support from the Singapore National Research Foundation through a fellowship
grant (NRF-RF2009-06), a tier2 grant from Singapore Ministry of
Education (MOE2011-T2-2-051), and Nanyang Technological
University for a start-up grant support (M58110061) and the
New Initiative Fund (M58110100). J. A. acknowledges the
funding from the Spanish MICINN projects MAT2010-15138
(COPEON) and CSD2009-00013 (IMAGINE) and Generalitat
deCatalunya (2009 SGR 770, NanoAraCat and XaRMAE). M. dlM
acknowledges the CSICJAE-Doc. The authors thank the TEM
facilities at the Universitat de Barcelona. We also acknowledge
Rosa Córdoba at LMA-INA for help in the FIB sample
preparation.
Pattern fabrication
Patterning on mica was achieved by laser interference lithography.42 Briey, a positive photoresist S1805 (MicroChem) was
spin-coated on a freshly cleaved thin layer of mica, followed by a
3 min baking step at 115 C. The mica substrate was then
exposed two times under two beams split from a 442 nm He–Cd
gas laser, with a 90 rotation of the substrate performed
between the two consecutive exposures. The formation of the
pattern was nalized by immersing the exposed substrate into
the corresponding developer MF-319 for 1 min and rinsing with
deionized water.
7248 | Nanoscale, 2013, 5, 7242–7249
Notes and references
1 J. E. Ayers, Heteroepitaxy of Semiconductors: Theory, Growth,
and Characterization, CRC Press, 2007.
2 J. W. Matthews and A. E. Blakeslee, J. Cryst. Growth, 1974, 27,
118–125.
3 F. Glas, Phys. Rev. B: Condens. Matter Mater. Phys., 2006, 74,
121302.
4 E. Ertekin, P. A. Greaney, D. C. Chrzan and T. D. Sands,
J. Appl. Phys., 2005, 97, 114325.
5 A. Koma, Thin Solid Films, 1992, 216, 72–76.
6 A. Koma, J. Cryst. Growth, 1999, 201–202, 236–241.
7 A. Koma, K. Sunouchi and T. Miyajima, Microelectron. Eng.,
1984, 2, 129–136.
8 A. Koma and K. Yoshimura, Surf. Sci., 1986, 174, 556–560.
This journal is ª The Royal Society of Chemistry 2013
View Article Online
Published on 13 May 2013. Downloaded on 31/08/2013 08:32:10.
Paper
9 K. Saiki, K. Ueno, T. Shimada and A. Koma, J. Cryst. Growth,
1989, 95, 603–606.
10 M. I. B. Utama, Z. Peng, R. Chen, B. Peng, X. Xu, Y. Dong,
L. M. Wong, S. Wang, H. Sun and Q. H. Xiong, Nano Lett.,
2010, 11, 3051–3057.
11 M. I. B. Utama, F. J. Belarre, C. Magen, B. Peng, J. Arbiol and
Q. H. Xiong, Nano Lett., 2012, 12, 2146–2152.
12 J. Pan, M. I. B. Utama, Q. Zhang, X. F. Liu, B. Peng,
L. M. Wong, T. C. Sum, S. J. Wang and Q. H. Xiong, Adv.
Mater., 2012, 24, 4151–4156.
13 M. I. B. Utama, Q. Zhang, S. F. Jia, D. H. Li, J. B. Wang and
Q. H. Xiong, ACS Nano, 2012, 6, 2281–2288.
14 Y. J. Hong, W. H. Lee, Y. Wu, R. S. Ruoff and T. Fukui, Nano
Lett., 2012, 12, 1431–1436.
15 H. Peng, W. Dang, J. Cao, Y. Chen, D. Wu, W. Zheng, H. Li,
Z. X. Shen and Z. Liu, Nat. Chem., 2012, 4, 281–286.
16 R. Viswanathan, J. A. Zasadzinski and D. K. Schwartz,
Science, 1993, 261, 449–452.
17 M. G. Astles, Liquid-Phase Epitaxial Growth of III–V Compound
Semiconductor Materials and Their Device Applications, Taylor
& Francis, 1990.
18 H. Q. Le, S. J. Chua, K. P. Loh, E. A. Fitzgerald and Y. W. Koh,
Nanotechnology, 2006, 17, 483–488.
19 B. Liu and E. S. Aydil, J. Am. Chem. Soc., 2009, 131, 3985–
3990.
20 I. Levin, A. Davydov, B. Nikoobakht, N. Sanford and
P. Mogilevsky, Appl. Phys. Lett., 2005, 87, 103110.
21 J. Shi and X. Wang, J. Phys. Chem. C, 2010, 114, 2082–2088.
22 M. H. Grabow and G. H. Gilmer, Surf. Sci., 1988, 194, 333–346.
23 N. A. Denesyuk and J.-P. Hansen, J. Chem. Phys., 2004, 121,
3613–3624.
24 C. Ratsch and J. A. Venables, J. Vac. Sci. Technol., A, 2003, 21,
S96–S109.
25 K. Govender, D. S. Boyle, P. B. Kenway and P. O'Brien,
J. Mater. Chem., 2004, 14, 2575–2591.
This journal is ª The Royal Society of Chemistry 2013
Nanoscale
26 J. G. E. Gardeniers, Z. M. Rittersma and G. J. Burger, J. Appl.
Phys., 1998, 83, 7844–7854.
27 D. Xia, Z. Ku, S. C. Lee and S. R. J. Brueck, Adv. Mater., 2011,
23, 147–179.
28 P. Yang, H. Yan, S. Mao, R. Russo, J. Johnson, R. Saykally,
N. Morris, J. Pham, R. He and H.-J. Choi, Adv. Funct.
Mater., 2002, 12, 323–331.
29 L. E. Greene, M. Law, J. Goldberger, F. Kim, J. C. Johnson,
Y. Zhang, R. J. Saykally and P. Yang, Angew. Chem., Int. Ed.,
2003, 42, 3031–3034.
30 Z. L. Wang, J. Phys.: Condens. Matter, 2004, 16, R829–R858.
31 C. Klingshirn, Semiconductor Optics, Springer, 2007.
32 Y. P. Varshni, Physica, 1967, 34, 149–154.
33 Z. R. Tian, J. A. Voigt, J. Liu, B. Mckenzie, M. J. Mcdermott,
M. A. Rodriguez, H. Konishi and H. Xu, Nat. Mater., 2003,
2, 821–826.
34 J. Joo, B. Y. Chow, M. Prakash, E. S. Boyden and
J. M. Jacobson, Nat. Mater., 2011, 10, 596–601.
35 O. Lupan, T. Pauporte, T. Le Bahers, B. Viana and I. Cioni,
Adv. Funct. Mater., 2011, 21, 3564–3572.
36 T. T. Hang, T. X. Anh and P. T. Huy, J. Phys.: Conf. Ser., 2009,
187, 012022.
37 K. Ueno, T. Shimada, K. Saiki and A. Koma, Appl. Phys. Lett.,
1990, 56, 327–329.
38 K. Y. Liu, K. Ueno, Y. Fujikawa, K. Saiki and A. Koma, Jpn. J.
Appl. Phys., 1993, 32, L434–L437.
39 S. Xu, C. Xu, Y. Liu, Y. Hu, R. Yang, Q. Yang, J. H. Ryou,
H. J. Kim, Z. Lochner, S. Choi, R. Dupuis and Z. L. Wang,
Adv. Mater., 2010, 22, 4749–4753.
40 S. Chu, G. Wang, W. Zhou, Y. Lin, L. Chernyak, J. Zhao,
J. Kong, L. Li, J. Ren and J. Liu, Nat. Nanotechnol., 2011, 6,
506–510.
41 L. Vayssieres, Adv. Mater., 2003, 15, 464–466.
42 Y. Wei, W. Wu, R. Guo, D. Yuan, S. Das and Z. L. Wang, Nano
Lett., 2010, 10, 3414–3419.
Nanoscale, 2013, 5, 7242–7249 | 7249