Materials Science and Engineering R 43 (2004) 103–138
Nanometric superlattices: non-lithographic fabrication,
materials, and prospects
H. Chik, J.M. Xu*
Division of Engineering, Brown University, Box D, Providence, RI 02912, USA
Abstract
A non-lithographic technique that utilizes the self-organized, highly ordered anodized aluminum oxide (AAO)
porous membrane as a template is employed as a general fabrication means for the formation of vastly different twodimensional lateral nanometric superlattices. The fact that material systems as different as metals, semiconductors,
and carbon nanotubes (CNT) can be treated with the same ease attests to the generality of this nano-fabrication
approach. The original alumina nanopore membranes determine the uniformity, packing density, and size of the
nanostructures. The flexibility of using a variety of materials, the accurate control over fabrication process, and the
command over the alumina template attributes give us the freedom of engineering various physical properties
determined by the shape, size, composition, and doping of the nanostructures. The novel nanomaterial platform
realized by this unique technique is powerfully enabling for a broad range of applications as well as for uncovering
new physical phenomena such as the collective behavior of arrays of nano-elements that may not be intrinsic to
individual nano-elements.
# 2003 Elsevier B.V. All rights reserved.
Keywords: Superlattices; Non-lithographic; Collective behavior; Nanotechnology; Quantum dot; Semiconductor devices
1. Introduction
1.1. Motivation: a single’s nano world thus far
Nanotechnology has become the subjects of much interest, intensive research, and in some
instances, a profitable enterprise. However, the march towards nanotechnology has been underway
long before it became fashionable, and in this regard, it is merely a natural step along the onedimensional path of miniaturization in the microelectronics evolution. Nevertheless, the ability to
engineer at the nano-scale does offer us something new and broadly important. It offers us access to a
qualitatively new paradigm in science and engineering, largely because the nano-scale is comparable
to the ‘‘electron wavelength’’ in electronic materials and to the sizes of most biological molecules.
From this perspective, engineering materials on the nano-scale is more than a fashion and worth of
investing in, for it can and will enable controlled alteration, dynamical manipulation, and molecular
functionalization of material properties, and potentially create entirely new properties that are not
intrinsic to the material in the bulk and in individual nano-element forms and open access to regimes
of material functions previously inaccessible.
*
Corresponding author.
E-mail address: [email protected] (J.M. Xu).
0927-796X/$ – see front matter # 2003 Elsevier B.V. All rights reserved.
doi:10.1016/j.mser.2003.12.001
104
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Technologically, the pursuit for nano-scale engineering has been powered by modern
lithographic and contact printing techniques, and has been mostly focused on explorations of
individually crafted nanostructures. Many of these achievements are astonishing including:
single nanotube field-effect transistor [1];
single electron transistor [2];
logic functionality from nano-networks [3,4];
single nanowire electrically pumped laser [5], etc.
As a first step into the nano-realm, investigation of isolated nanostructures and characterization
of single nanodevices are of great importance and naturally of strong appeal to researchers.
However, a more challenging and perhaps of great potential for novel applications is the
collective behavior of nanostructures in a large ensemble and the manipulation of multiple
nanostructures into integrable functional units during which the properties of the system as a whole
may be fundamentally altered and have no classical or individual-based characteristics. Thus, it is
both desirable and natural to take the next step and seek for ways to advance both the fabrication
capabilities and the science explorations beyond the realm of individual nanostructures.
1.2. Collective behavior
A not yet heated but no less interesting space of nanostructured materials is that of the collective
properties (behaviors) of large assembles of coupled nano-elements giving rise to emergent
responses of a system as a whole that are not intrinsic to individual nano-elements. Collective
behavior is more complex than examining the properties of the elements in isolation as it involves
unique interactions of a group of nano-elements acting together, however, a better understanding of
the intricacies can be highly rewarding in the design principles of robust, flexible, and decentralized
physical systems that are resistant to errors, adaptive, self-organize, and capable of acquiring
information and executing on the information. One example of such a system is the ant colonies as
the communication and interaction between relatively simple individual organisms give rise to an
effective survival strategy for the colony as a whole. Borrowing from biology, a more complex
example would be the behavior of individual neurons and how they act together in the human brain
to perform bodily functions.
In photonics, one extraordinarily well popularized example of interesting behaviors arising from
couplings is the so-called photonic band gap crystals, which is essentially an extension of the Bragg
gratings from one dimension to three dimensions in real space, that rely on couplings on the length
scale of half or quarter optical wavelengths. In the near-field coupling regime, the spacing between
adjacent nano-elements is on the nano-scale, much smaller than the optical wavelengths. As such, it
enables drastically different coupling mechanisms, both the short-range couplings via electronic
tunneling and the short-range photonic tunneling (that is, coupling via an optical field not yet
emitted), as well as the long-range EM field couplings. In this regime, synchronized (coherent)
resonant excitations and responses of the nano-elements can be extraordinary and give rise to
unexpected or entirely new properties or behaviors in otherwise familiar and ordinary materials.
Understanding of the unique properties arising from the collective behaviors of such lateral
superlattices in the presence of near-field couplings due to the sub-wavelength and periodic spacing,
and coherent interactions with an incident or emitted radiation field will enable us to explore, design,
and demonstrate novel nano-electronic, nano-photonic, and nano-magnetic devices, and functional
subsystems based on this new class of nanostructured materials.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
To begin to investigate the interesting regime of strong collective electronic and photonic
couplings, structures and systems with critical feature sizes much smaller than the wavelength of
interest need to be fabricated. However, the features of the desired nanometric structures (2D or 3D
superlattices) including good nano-element uniformity over a large scale area are often beyond the
reach of conventional fabrication techniques.
1.3. Scope and aim of the paper
It is intuitively clear that maximum collective resonance effects in this regime would require
large arrays of highly ordered and highly uniform nano-elements. Fabricating such large arrays of
nanostructures poses a challenge to even the best developed nano-fabrication technique—the
electron beam lithography, because of its low throughput due to its long exposure time, small field
size, and high cost of equipment. On the other hand, a few new nano-fabrication approaches appear
to be well poised to meet this challenge. A non-lithographic nanopore template based approach is
particularly encouraging and is the chosen focus for this review.
The non-lithographic nano-fabrication approach of interest to us in this article utilizes highly
ordered hexagonally packed straight nanopore arrays formed via self-organization in a carefully
controlled process of anodizing high-purity aluminum. Using the anodized aluminum oxide (AAO)
membrane as a template, a set of non-lithographic methods for nano-fabrication have been
developed and applied to the formation of a wide variety of nanometric superlattices including
nanodots, nanowires, nanopillars, and nanopores. The fabrication of these nanostructures have been
demonstrated in metals, semiconductors, and polymers, and will be described in this article. Taking
advantage of its hexagonal symmetry, the highest packing form for a given diameter and spacing,
applications utilizing the periodic nanometric superlattices will also be described.
2. AAO membrane: a highly ordered template
2.1. Properties
Fabricating nanostructures and tailoring their properties on the length scales less than 50 nm by
conventional lithographic approaches is quite challenging. And, making them into large highly
ordered and uniform arrays, to enable and enhance desirable collective behaviors, is even more
difficult. Such difficulties have compelled many including us to seek alternative approaches such as
utilization of recent developments in chemical synthetic methods to produce uniform nanostructures
on these length scales, and to form self-organized highly ordered superstructures (i.e. multidimensional superlattices).
Central to the non-lithographic, nano-fabrication method favored here is the utilization of an
alumina nanopore membrane as either the growth template, or the etch or evaporation mask for the
formation of the proposed nanostructures. The alumina nanopore array is formed by anodization of
high-purity aluminum under certain carefully controlled anodization conditions. The nanopores can
self-organize into a highly ordered array of uniform pores with the pore diameter, the period, and the
array size being variable over ranges that are beyond the reach of standard e-beam lithography. As an
example, Fig. 1 is a hexagonally packed template with 55 nm diameter pore sizes spaced 110 nm
apart.
Anodic porous alumina has been studied in detail over the last five decades [6–9]. In the
anodization process, an electrical circuit is established between a cathode and a thin film of
105
106
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 1. Anodized aluminum oxide (AAO) membrane with pore diameter of 55 nm and pore spacing of 110 nm: (a) top
view and (b) cross-sectional view.
aluminum which serves as the anode, according to following reaction:
2Al þ 3H2 O ¼ Al2 O3 þ 3H2 ;
DG ¼ 864:6 kJ
where DG is the standard Gibbs free energy change. During the anodization, initially a planar
barrier film forms followed by pore development leading to the formation of the relatively regular
porous anodic film, which thickens in time [10]. This barrier film is maintained during further oxide
growth as a semi-spherical oxide layer at pore bottoms as shown in the schematic in Fig. 2. The
formation of regular nanopore structures actually means that, for some reason, micro-heterogeneity
of the oxide comes into play resulting in easier dissolution of the oxide at some points than at others,
thus implying a varying dissolution rate constant over the surface. A steady-state pore growth regime
is reached characterized by the balance between the field-enhanced oxide dissolution at the oxide/
electrolyte interface at the pore bottoms and the formation of oxide at the metal/oxide interface due
to migration of O2/OH ions through the bottom oxide layer. The cell and pore diameters, and the
barrier layer thickness are directly proportional to the formation voltage [11].
In 1988, Konno et al. [12] observed pore regularity in anodized aluminum oxide membranes. By
1995, Masuda and Fukuda [13] reported large defect-free regions appearing in large domains, while
defects were found at the boundaries of these domains after a long anodization period under an
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 2. Schematic of the semi-spherical oxide layer at the bottom of the AAO nanopore membrane.
appropriate constant anodizing voltage. The long anodization time rearranges the cells and reduces
the number of defects and dislocations. Bandyopadhyay et al. [14] showed that pores tend to
nucleate in regions of increased surface elastic energy and grain boundaries are typically preferred
sites. Although there is no direct evidence up to now that defects in the pore lattice can originate
from dislocations or other defects in the aluminum, it is clear that the correlation of self-organization
is at least disturbed by a large number of grain boundaries in the non-annealed aluminum foils. The
high surface roughness leads to a faster formation of barrier oxide and pores at depressions in the
surface compared to other smoother locations. The pores that nucleated at such depressions in an
early stage will grow sooner than at other sites. The surface roughness is thus transferred to the
anodization front, at the interface between the aluminum and oxide layer, and prevent selforganization. However, the stirring of the electrolyte may contribute to spatially homogenous etching
conditions. Pre-annealing of the aluminum substrate can increase the grain size, while stirring of the
electrolyte during oxidation, and using smooth electropolished aluminum surfaces are all necessary
for obtaining ordered hexagonal structures with typical long-range ordering of 100 mm2.
A two-step anodization process, illustrated in Fig. 3, can improve the pore regularity under
particular conditions. In this process, after stripping away the thick aluminum oxide film obtained
from the first long anodization, a porous thin alumina film with highly ordered pores can be obtained
by a subsequent re-anodization under the appropriate conditions.
While the physical mechanisms underlying this remarkable self-organization process are not yet
understood, some compelling but incomplete hypothesis and models have been advanced in the
literature. For example, a mechanical stress mechanism has been proposed to explain the selfordering of anodic alumina [15]. It was suggested that the repulsive forces between neighboring
pores caused by mechanical stress at the metal/oxide interface promote the formation of hexagonally
ordered pore arrangements. According to this hypothesis, under the usual experimental conditions,
the volume expansion accompanying the aluminum oxidation gives rise to the mechanical stress at
Fig. 3. Schematic of the two-step anodization process. The first anodization step pretextures the Al surface in a highly
ordered manner which is retained during the second anodization.
107
108
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
the interface between alumina and Al. The expansion strongly depends on experimental conditions
such as the electrolyte concentration and anodizing voltage. The optimal conditions for the growth of
ordered arrangements are accomplished by moderate expansion of the aluminum, whereas no
ordered domains can be observed in the cases of contraction or very strong volume expansion. These
observations suggest that the system could be modeled by a Gibbs energy function consisting of an
elastic energy term, a volume energy term, and a surface energy term. The complication comes from
the role of the chemical energy which has remained elusive from definitive experimental
characterization, or explicit functional expression of key parameters.
After the second-step anodization, the nanopore alumina can be separated from the Al base and
further processed into a freestanding membrane of nanopores that is open on the top and bottom and
be used as a base template stencil or mask for fabricating a variety of highly ordered nanostructures.
The advantages of the non-lithographic fabrication using self-organized, highly ordered anodized
aluminum oxide nanopore template include:
uniform pore diameter adjustable from 20 to 200 nm;
uniform pore periodicity from 50 to 400 nm;
highest packing density 109–1011 cm2 due to its hexagonal symmetry;
the pore diameter, the period, and the array size variable over ranges that are beyond the reach of
standard e-beam lithography with the AAO template pattern completely replicated in
dimensionality onto the receiving substrate;
large area and low cost;
process is not material specific ranging from oxides to semiconductors to metals to polymers;
readily scalable and compatible with existing IC processing yet inherently not limited to existing
wafer sizes;
process conformal to arbitrary surfaces either flat or curved surfaces, both locally or globally;
does not need a cleanroom process.
2.2. Pattern transfer
Periodic porous oxide with remarkably ordered arrangement of cylindrical pores whose
diameter variable in the range of 25–200 nm with pore spacings from 50 to 400 nm can now be
realized routinely. The process typical begins with a high-purity aluminum film (99.98%),
followed by electropolishing in mixed solution of HClO4 and C2H5OH for 3 min, and completed by
a two-step anodization, in either oxalic acid, sulphuric acid, or phosphoric acid solution over several
hours. Li et al. established that by varying the anodization conditions, hexagonal close-packed arrays
with controllable variation of diameters, densities, and lengths can be formed defect-free over large
areas [16]. A set of optimized anodization parameters in different electrolytes are shown in Table 1.
The resultant nano-scale pores are parallel, straight, and oriented vertical to the surface. In addition,
the nanopores are uniform in diameter and in spacing.
Table 1
Typical anodization conditions for various electrolytes
Electrolyte
Concentration (M)
Temperature (8C)
Voltage (V)
Typical pore diameter (nm)
Sulphuric acid
Oxalic acid
Phosphoric acid
0.5
0.3
1.0
0
10
0
25
40
160
30
45
400
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 4. Generalized schematic illustrating the process flow for various nanostructures.
After reaching the desired template thickness during the second anodization, the AAO
membrane can be separated from the remaining aluminum foil using a 2% HgCl2 solution or by
reversing the applied voltage of anodization to a negative bias. The resultant nanopores are open on
one side of the membrane, while a thin Al2O3 barrier layer exists on the opposite side. Both wet
etching in a 0.1 M phosphoric acid solution and plasma etching with CF4 þ O2 gas system are
effective for removing the barrier layer and forming a through-pore membrane [19].
By using the AAO membrane as templates, stencils, or masks, a number of approaches for
fabricating a variety of nano-superstructures have been developed. These approaches are not
material specific. They are also scalable in principle, however, in practice the sizes of the arrayed
nanopore membranes have been limited to the size on the order of a couple of centimeters due to the
fact that they are sufficient for most research and development purposes and that large membranes
are brittle to handle manually in the lab. The process flow diagram in Fig. 4 illustrates some of the
possible ways of creating nanostructured arrays. In the sections to follow, we will describe the
details of the fabrication processes for each of these primary classes of nanometric superlattices,
namely:
1.
2.
3.
4.
nanodot;
nanowire and nanopillar;
nanotubes;
nanopores or nano anti-dots.
It is already clear from the different pathways outlined in the schematic that the AAO templating
method is extremely versatile and can be applied to a broad range of purposes and devices.
3. Nanodots
3.1. Fabrication
The placement of metal nanodots on a substrate in a controlled manner has been achievable with
the advent of e-beam lithography. Current e-beam writing machines can readily write one or a
hundred or even thousands of nanodot patterns in e-beam photoresist with great accuracy and
109
110
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
without extraneous efforts. However, when asked to write hundreds of millions of nanodots,
the process would take extraordinary amount of time and also be limited by the microscopes field
of view if stitching errors are to be avoided. Also, this does not account for complex proximity
effect issues when deciding on exposure times which can dramatically alter the dimensions of the
desired patterns especially with nanodots spaced about 100 nm or less. With the AAO templating
method, billions of nanodots of uniform size and spacing can be formed in parallel often in one or a
few steps.
The fabrication of metal nanodots starts with the nanopore AAO membrane with the barrier
layer removed leaving a through-pore template. The template is then placed onto the substrate done
usually in solution as the membrane, being as thin and as large as possible, can be easily fractured in
manual handling. The template adheres well and strongly to a flat substrate because of the Van der
Waals force. Afterwards, a thin metal layer can be evaporated through the pores onto the substrate
surface. Evaporation is typically performed in an e-beam evaporator at a relatively slow deposition
rate as faster rates tend to clog the nanopores more easily. After peeling off or etching away the AAO
template, undesired metal is also lifted off, much like a conventional lift-off process in
semiconductor processing, and a perfect hexagonal array of metal nanodots is formed naturally
on the hosting (e.g. semiconductor) substrate. The mean diameter and spacing of the nanodots are
limited by, and thus can be varied with, the pore diameter and spacing of the AAO membrane. The
dot diameter can be reduced further post lift-off by wet chemical etching. A typical nanodot array is
shown in Fig. 5. Here, a Ni nanodot array on a Si substrate with 55 nm diameter nanodots and a
periodicity measuring the center-to-center spacing between two neighboring dots 110 nm is shown in
the SEM image. The height of the metal nanodots, determined to be 30 nm by AFM, can be adjusted
by controlling the evaporation thickness or by post-evaporation etching. To reach a metal nanodot
height of 30 nm, 100 nm thick Ni was evaporated as indicated by the thickness monitor. The
nanodots are in the shape of a cone likely due to the finite solid angle of the impinging beam and due
to the relatively large aspect ratio of the pore (50:500 nm). It is also possible that as evaporation
continues, metal accumulates at the bottom of the nanochannel (surface of substrate), but also on the
sidewalls reducing the diameter of the nanochannel resulting in a conical shaped nanodot. The highly
ordered magnetic (Ni) nanodot arrays provide us an enabling platform for exploiting the collective
Fig. 5. Ni nanodot array with 55 nm diameter and 110 nm spacing mimicking the AAO template dimensions. The scale bar
is 100 nm.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 6. MBE grown GaAs nanodot array on GaAs substrate with (a) top view and (b) oblique angle view from [20].
behaviors of such arrays for low-power, room temperature, high-density information storage, as well
as information processing.
The process is non-selective to either target metal materials or substrate materials. Different
metal nanodot arrays, such as Au [17], Ni [18,19], Co, and Fe, have been deposited in a similar
fashion on various substrates such as Si, GaAs, and GaN. The ability to form metallic nanodots on a
substrate provides the platform for etching nanopillars and growing nanowires, as will be described
in proceeding sections. However, nanodots need not to be confined to just the metallic variety
although evaporating metal through the nanopore template is a simple process, GaAs nanodot arrays
have been grown on GaAs (0 0 1) substrates through the nanopore template by molecular beam
epitaxy (MBE) [20]. Mei et al. observed disk-like shaped with a slightly convex surface as shown in
Fig. 6, but also pyramidal-like GaAs nanodots depending on the growth rate used. Higher growth
rates produced disk-like nanodot shapes indicating layer-by-layer growth mode, however, greater
portions of the GaAs material deposited on the AAO template surface and not within the
nanochannels on the surface of the substrate. The ratio of material depositing inside the nanochannel
or on top of the nanopore membrane was improved with lower growth rates, however, this resulted in
the pyramid-like shaped nanodots indicating the growth mode was not layer-by-layer. The thickness
of the AAO template was from 120 to 150 nm allowing the Ga beam to penetrate through
the nanochannels as the growth mechanism is migration dominated with no nanodot growth
when templates are greater than 300 nm thick. From these demonstrations, we observe that a third
111
112
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 7. Two spatially resolved Au nanodots evaporated through each nanopore on Si from [21].
deposition route seems to be feasible as well in which a metallic nanodot array, such as Cd, is formed
by evaporation then converted into compound semiconductor nanodots such as CdS.
Beyond the capabilities of e-beam lithography is the formation of nanodot arrays whereby each
nanodot comprises of multiple metals in a 3D structure. Fabrication of ordered arrays of multimaterial composite nanodots maybe desirable to enable new optical or catalytic properties of
nanoparticles. Using the AAO template as an evaporation mask, identical to the single metal nanodot
formation as described previously, and by tilting the sample (AAO template attached onto the
substrate) to an appropriate angle, shadowed evaporation can occur. Shadowed evaporation creates
spatially resolved deposition as the effective aperture, i.e. metal with a straight line path from the
template opening to the bottom of the nanochannel, diminishes to only one section of the nanopore at
the surface of the substrate. Masuda et al. deposited two and three nanodots with the same or differing
composition in each nanopore as reported in Fig. 7 [21]. In the SEM image shown, two Au nanodots
sit on a Si substrate produced by a single nanochannel. A bimetallic nanodot array can be created
simply by switching the metals being evaporated. The authors demonstrated Au and Ag nanodots as
well as three (Au, Ag, and Si) nanodots created by a single nanochannel.
3.2. Examples of applications
Circular nano-magnet arrays are being developed for use in random access memory (RAM)
devices. Magnetic nanodot arrays also can be deployed to affect Faraday rotation in waveguides or
vertical cavity surface emitting lasers (VCSELs). The template method allows large-scale fabrication
of uniformly sized and site controlled magnetic nanodot arrays, and offers a number of advantages,
such as smaller minimum sizes over conventional photolithography, and larger array size and lower
costs over e-beam lithography.
Two of the biggest problems today’s microelectronics face are power dissipation and
interconnection. When compared to the also significant challenges of device miniaturization and the
major costs associated with building next generation fabrication lines, these two problems are more
fundamental and rooted in the binary serial architecture. While nano-scale devices based on
molecular-scale elements have been demonstrated, such devices still face the difficult interconnection problems if large-scale integration is to be attempted. Furthermore, since power dissipation from
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
prototype molecular devices is on the order of microwatts, thermal management will continue to be a
major design challenge even if devices are greatly reduced in size [22,23].
The study of magnetic nanostructures could enable us to move beyond the realm of information
storages into information processing. Nanocomposite materials and engineered nano-magnetic arrays
appear to be promising systems for exploring magnetic interaction modes which are not possible in
naturally occurring or bulk materials. Additionally, nano-magnetic dots not only have the potential to
each represent 1 bit of information, which can be highly stable, but also could act collectively to
process and transmit information in a cellular automata or a logic circuitry configuration at room
temperature. Nano-scale logic using magnetic nanostructures could have considerable value for
computation and signal processing. Potential advantages of such systems include low-power
dissipation and low interconnection demand, thermal stability at room temperature, and high levels of
integrability and scalability [24].
4. Nanowires and nanopillars
4.1. Metal nanowires
Arrays of nano metallic and magnetic wires are attractive for their potential applications in
radiation and magnetic sensing, high-density magnetic devices, surface enhanced Raman spectroscopy, as well as for fundamental studies of nano-magnetics. The ability to produce highly ordered
nanowire arrays cheaply and effectively is important for both purposes. AAO template synthesis
using electrodeposition has found to be stable at high temperatures and in organic solvents. It has
proved to be a low cost, high yield, and high throughput technique for producing large arrays of
nanowires.
The primary advantage of this method is ease of material handling. The aluminum substrate
remains to supply both mechanical support and electronic contact. The thickness of the oxide layer
can be varied freely from very thin to very thick. Various metals such as Ag [25], Fe [26,27], Ni
[27,28], and Bi [29] just to name a few have been deposited into the nanochannels by an alternating
current (ac) deposition technique. Only ac conditions is found to be effective due to the restrictions
arising from the barrier layer of the AAO template and also due to diffusion in the nanopores. An
example application of metallic nanowires is uncooled IR detection operating in bolometer mode.
Bismuth nanowires are expected to manifest strong quantum size effects due to the large Fermi
velocity in Bi and the small transverse dimension. The GMR effect in the magnetic nanowire
configuration is expected to exhibit a critical dependence on both the wire diameter and the spacing
as well as a great increase in magnitude (possibly even in sign) due to the 1D density of states.
Many of the electrodeposition processes have exhibited the so-called ‘‘sky-scraper’’
phenomenon referring to the lack of length uniformity and control. Common to these processes is
the use of aqueous solutions, with the exception of a few that used organic solutions [30–32]. In
these processes, the positive feedback makes taller wires grow faster and taller. Other drawbacks
associated with electrodeposition in aqueous solution include the relatively limited number of
elements that can be deposited from aqueous electrolytes, and the need for a conducting substrate.
Additionally, the quality of the nanowire arrays may be affected by a number of side reactions, such
as oxidation of the metal, and water electrolysis. Eventually, stoppage of wire growth may occur due
to the sealing of the AAO template nanopores.
By changing the electrolyte of the electodeposition process into a dimethylsulphoxide (DMSO)
solution containing metal chloride, the drawbacks associated with aqueous bath can be overcome.
113
114
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 8. The effect of ac frequency on the filling percentage of the nanopores for Ni nanowire arrays: (a) 0.1 kHz,
(b) 0.3 kHz, (c) 0.5 kHz, and (d) 0.75 kHz.
For example, NiCl2 and BiCl2 in DMSO was used to grow Ni and Bi nanowires, respectively [33].
With an appropriate choice of the electrodeposition parameters, such as the electolyte concentration,
the temperature, and the bath voltage, it is possible to control the growth rate of the nanowire arrays
with a remarkable degree of precision and uniformity. In addition, the ac bias parameters, in
particular the frequency and waveform affects the filling ratio into the nanopores as shown in Fig. 8.
Higher frequencies result in near complete filling of the nanopores as more crystalline nuclei forms
resulting in easier deposition of the metal into the nanopores and promoting homogeneous growth of
the nanowires. In Fig. 9, the Ni and Bi nanowires were partially exposed by etching back the AAO
template allowing easy visualization of the array uniformity. As seen in the image, the nanowires
form a highly ordered hexagonal array of uniform spacing and diameter with a good level of
uniformity and control in the length. The significance of these results can easily be appreciated in the
field of computing technology.
Since the early 1990s with the gaining popularity of the personal computer, the storage density
of hard disks have increased about 60% per year. In order to ensure good signal-to-noise ratios, each
bit on the recording media must be composed of many grains (1000), thus as bit sizes are reduced
to increase the areal density, the individual grains are reduced accordingly to a point where thermal
stability will become an issue. Perpendicular recording media instead of the traditional longitudinal
recording method has attracted great research interest due to the belief that it will offer the largest
storage densities. Here, 1 bit corresponds to one single domain particle or a nano-magnet. A highly
ordered array of nanowires is a natural choice with each nanowire representing a single nanomagnet. Because of their nano-scale size, anisotropic shape, and spatial separation from each other,
each nanowire is a single domain with only two perpendicular magnetization states (i.e. up or down).
With a lattice constant of about 50 nm, an areal density of about 300 Gbit/in.2 can be achieved by the
hexagonally arranged array. In 1994, Chou et al. demonstrated the ultrahigh density storage using
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 9. Ni nanowires (a) top view and (b) AAO partially etched exposing the nanowires. Bi nanowires (c) after AAO was
dissolved, and (d) cross-sectional view in which the bright lines are the nanowires exposed by cracking the AAO template.
single domain magnetic nanopillar array by lithography [34]. The 35 nm diameter Ni nanowires with
a period of 100 nm were fabricated using advanced electron beam lithography and stamping and
electroplating techniques.
Using the much less demanding AAO template-assisted nanowire fabrication technique, Nieisch
et al. created a similar Ni nanowire array [35]. With a nanowire diameter of 30 nm, the coercivity
was 1200 Oe and a squareness of 98% when the external field was applied parallel to the length of
the nanowire as shown in Fig. 10(c). Each Ni nanowire is expected to be able to be switched
independently to the magnetization of its nearest neighbors and thus be able to store 1 bit of
information as shown in the magnetic force microscopy image in Fig. 10(b). The magnetic properties
of Fe nanowire arrays prepared by the AAO templating method have also been investigated [27].
An alternative method to grow metallic nanowires involves first removing the barrier layer of
the AAO membrane followed by evaporating a metal layer onto one surface of the through-pore
template providing a contact for electrodeposition. Various metals can then be deposited into the
nanochannels by a simple direct current (dc) deposition technique. The drawback of this technique is
the requirement of a freestanding AAO template that must be mechanically robust enough to tolerate
manipulation since the membrane is already separated from the supporting bulk aluminum substrate.
On the other hand, a thicker template makes it more difficult to fill each nanopore in an uniform
fashion.
4.2. Carbon nanotubes
Since its discovery in 1991 [36], carbon nanotubes (CNT) has invoked tremendous research
interests as they appear to be amongst the most promising materials anticipated to impact future
nanotechnology. Their unique structural and electronic properties [37,38], found to vary from metallic
115
116
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 10. Ni nanowires grown in AAO template from [35]: (a) top view; (b) MFM image showing the nanowires magnetized
in the ‘‘up’’ (light) and ‘‘down’’ (dark) states, and (c) SQUID hysteresis loops for Ni nanowires with a diameter of 30 nm
and a spacing of 100 nm.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 11. Process flow for CVD carbon nanotube growth in AAO template: (I) formation of AAO template, (II)
electrodeposition of catalyst, (III) CVD growth, and (IV) post growth processing to expose the nanotubes.
to semiconducting depending on their lattice symmetry and dimensions, have envisioned for
prospective applications include field emission displays, biocompatible sensors, and molecular scale
electronic and photonic devices to name just a few. Most of these applications will require a
fabrication method capable of producing uniform carbon nanotubes with well-defined and
controllable properties. Additionally, the technique employed must be reproducible if any largescale fabrication of carbon nanotubes are envisioned. Efforts to fabricate aligned, isolated carbon
nanotubes in an array form have lead to some degrees of success [39–43], arrays of parallel carbon
nanotubes with uniform diameters and periodic arrangements have only been achieved with the use
of the AAO template [44–48].
The fabrication procedure begins at the same starting point as before, namely after the second
anodization step where the AAO template consists of uniform, parallel, hexagonally packed
nanochannels following the fabrication procedure previously described. The next step is to
electrochemically deposit a small amount of cobalt, nickel, or iron catalyst into the bottom of the
template channels as illustrated in the schematic shown in Fig. 11. The ordered array of nanotubes
are grown by first reducing the catalysts by heating the cobalt-loaded templates in a tube furnace at
600 8C for 4–5 h under a CO flow at a rate of 100 sccm. The CO flow is then replaced by a mixture
of 10% acetylene in nitrogen at the same flow rate for 2 h at 650 8C or higher. The samples are then
annealed in nitrogen for 15 h at the same temperature. Fig. 12 shows a SEM image of carbon
nanotube arrays which have been ion-milled to remove residual amorphous carbon from the template
surface. The nanotubes can then be partially exposed by etching the AAO membrane using a mixture
of phosphoric and chromic acid. The resultant nanotubes are uniform in length, parallel to each
other, and perpendicular to the template. The diameter of the nanotubes shown is approximately
47 nm which is slightly larger than the original template diameter due to uniform widening of
the template nanochannels during the catalyst deposition process. The diameter distribution is only
5% of the mean diameter much narrower than results reported using other methods to synthesis
carbon nanotube arrays. The length of the nanotubes is equal to the thickness of the AAO membrane
in which they were grown for which in this case was 6 mm. A cross-section SEM image shown in
Fig. 13 of the nanotube array partially exposed from the AAO template shows the cobalt catalyst is
still at the base of the nanotubes separated from the aluminum substrate. A clearer picture of the
cobalt catalyst remaining in the base of a separated nanotube is captured in the TEM image shown in
the inset of the figure.
The use of CNTs for large scale nano-electronics has hinged on being able to position the CNTs
selectively and to integrate them in ultra high densities. With the carbon nanotubes now positioned in
the periodically ordered AAO template and densities approaching 1011 order of magnitude elements
per square centimeter, this obstacle appears to have been lowered. Choi et al. presented a 20 nm in
diameter, 40 nm in pitch tera-level density CNT transistor grown in the AAO templates [49]. E-beam
117
118
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 12. Carbon nanotubes grown in AAO template (a) oblique angle view with nanotubes exposed and (b) top view
showing 47 nm diameter nanotubes.
lithography was used to pattern the AAO membrane before nanotube synthesis to selectively grow
CNTs in certain nanopores and after nanotube synthesis to define the gate, drain, and source contacts
of the transistor. Each CNT is electrically attached by rows of metal contacts on the top and bottom
surfaces forming a matrix of electrodes. To address an individual vertical CNT, a bias must be
applied between the corresponding row of electrode on the top and bottom surfaces. In this
configuration, each individual CNT is uniquely addressable as shown in Fig. 14. Taking advantage of
the unique properties of the CNTs such as high thermal power dissipation, high electrical
conductivity, and ultra small power consumption, it is conceivable that tera-level density CNTs will
play a role in the realization of next generation nano-electronics.
The synthesis of connections between two or more different carbon nanotubes is an important
step in the development of carbon nanotube based electronic devices and circuits. However, this is
difficult to achieve using conventional methods to grow carbon nanotubes because the straight tube
structure cannot be controllably altered along its length. Various ideas for post-growth modifications
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 13. Cross-section SEM image showing the presence of the cobalt catalyst sitting at the base of the carbon nanotubes.
The inset is a TEM image of a single nanotube.
have been suggested [50], but these have been hard to implement and are prone to defects. The AAO
template method to grow uniform, highly ordered straight carbon nanotubes as described above, can
be extended to grow uniform, highly ordered Y-junction carbon nanotubes [51]. The growth of the
Y-junction CNTs follows the same recipe as growing straight CNTs of pyrolysis of acetylene with
cobalt catalysis. The alteration comes in the preparation of the AAO template. Because the pore
diameter is proportional to the anodization voltage, reducing the voltage by a factor of 21/2 results
in twice as many nanopores appearing in order to maintain the of the template, and nearly all
nanopores branched into two smaller diameter nanopores. Thus, by switching the voltage during
anodization, a larger nanopore (stem) will gradually evolve into two smaller diameter nanopores
(branches) as shown in the inset of Fig. 15 where the resulting template consists of parallel
Y-branched nanopores with stems 40 nm in diameter and branches 28 nm in diameter. In the crosssection SEM image, the arrow indicates the transition between the stems (upper 3 mm of template)
Fig. 14. Carbon nanotubes grown in patterned AAO nanopores after [49]: (a) oblique angle view of exposed nanotubes, (b)
schematic of the transistor architecture, and (c) top view of an array of transistors with each at the cross point of a drain and
a gate electrode.
119
120
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 15. Cross-section view of a Y-junction AAO template with the arrow pointing to the straight to Y-junction CNT
growth transition. The inset depicts a clear image of the transition region.
and the branches (lower 2 mm of template). A TEM image in Fig. 16 shows the Y-junction CNT after
being removed from the template. In this case, the stems are 90 nm in diameter while the branches
are 50 nm in diameter.
Electronic transport measurements show that the Y-junctions act as semiconductor heterojunctions as their behavior can be explained simply by considering the change in bandgap across the
junction caused by joining together different diameter nanotubes [52]. This structural configuration
naturally leads to a three-terminal nano-scale transistor by applying different voltages to each of the
Y-junction arms. This integrated transistor function will provide the nano-electronics community
with a new base material for the development of molecular-scale electronic devices. With the
Y-junction CNTs still in the AAO template and each one behaving as a transistor, logic gates can be
formed by selectively connecting two or more nanotubes in parallel from the array as illustrated
schematically in Fig. 17.
4.3. Semiconductor nanopillars
Instead of growing nanostructures by electrochemical deposition as for metallic nanowires or by
pyrolysis for carbon nanotubes, in this section, the nanostructures are fabricated by etching into
existing material and using the AAO template as an etching mask in the traditional sense. In standard
lithography, to etch a pillar from a bulk substrate, a mask, usually formed from metal or silicon
dioxide/nitride, is deposited to protect desired areas from the etching chemicals. Everywhere else left
unprotected, the material is etched away leaving the a pillar-like structure protruding from the
substrate. Here a similar approach is taken to fabricate highly ordered nanopillars or superlattice of
quantum dots described in the next section.
The fabrication procedure begins with forming a perfect hexagonal array of metal nanodots onto
the substrate, as described earlier, where the nanopillars are to be formed. Then a deep reactive ion
etching (RIE) process is employed to obtain high aspect ratio semiconductor nanopillars using the
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 16. Y-junction CNTs: (a) oblique angle view with the CNTs still in the AAO template and (b) TEM image of the Yjunction region.
metal nanodots as the etching mask. Fig. 18 shows the oblique SEM views of GaN and GaAs nanowire
arrays after their respective etches using the metal nanodot array as a mask [53]. An aspect ratio of
1:10 was achieved with a nearly vertical etch profile. The dry etching conditions are listed in Table 2.
4.4. Quantum dots
The idea to further confine carriers to a greater degree and ultimately reducing the
dimensionality to zero, a quantum dot, has been sought after in the development of electro-optic
121
122
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 17. Schematic of example logic gates formed by selectively connecting two or more Y-junction CNTs, which are
individually a transistor, in parallel.
devices. When the quantum dots act as the active layer of a semiconductor laser, theoretically the
device offers low threshold current density, good temperature characteristics, and broad wavelength
tunability. To obtain maximum quantum effect, dots must be of uniform size, shape, and
composition. The most commonly used method to fabricate quantum dots so far is the selfassembled epitaxial growth based on strain-induced Stranski-Kranstanov (SK) mode growth [54–56].
Fig. 18. Oblique angle views of (a) GaN nanopillars and (b) GaAs nanopillars.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
123
Table 2
Typical dry etching conditions for nanopillar formation
Substrate
Gas
Flow rate (sccm)
Pressure (mTorr)
Power (W)
GaAs
GaN
Nb
Si
BCl3
Cl2
CF4 þ O2
BCl3 þ Cl2
20
60
42 þ 2
5 þ 10
15
80
50
50
100
200
100
100
It is, however, material specific as the growth mode relies on the strain between the quantum dot and
the substrate material such as InAs/GaAs, thus cannot be used in the growth of homoepitaxy and
closely lattice matched heterostructures such as AlGaAs/GaAs. Also, the SK growth mode still lacks
the uniformity, the ordering, and the packing density achievable in the nanopore template approach.
One alternative strategy capable of forming complex structured QD arrays is to start with
epitaxially grown multi-layered structures such as multi-quantum wells, PN and heterojunction
structures, followed by the use of the template-assisted method of fabricating high aspect ratio
nanopillar array as described earlier. One major advantage of this method for QD fabrication over
self-assembly or self-assembly based strain engineering is that it allows for accurate multi-activelayer stacking. The other advantages of this method over self-assembly include the uniformity and
the accurate controllability of feature size and position.
A schematic is shown in Fig. 19 to illustrate this idea where a multi-quantum well wafer is
patterned into three vertically stacked lateral superlattices of quantum dots, which are quantum
mechanically coupled. Moreover, each column of coupled quantum dots can be contacted naturally
on the top with self-aligned metal dots. An anticipated problem of this approach is the rather large
surface recombination usually occurring on un-passivated etched surfaces. A solution might be in
epitaxial regrowth. This problem seems not as severe in such material systems as GaN where the
free-surface recombination velocity is low as demonstrated in [57]. A comparison of templateassisted etched 50 nm nanopillars in InGaN/GaN multi-quantum well to an unprocessed wafer
revealed similar optical properties suggesting that surface recombination at the etch exposed vertical
walls of the nanopillars is modest at best. When compared to e-beam written 100 nm nanopillars, the
two nanopillars revealed robust photoluminescence (PL) efficiencies, not impaired by the very large
surface to volume ratio and even exceeding those of an unpatterned quantum well wafer when taken
into account the actual areal filling factors. This has presumably benefited from the higher photon
escape probability due to the post-like texture of the light emitting medium.
Fig. 19. Schematic of transforming an epitaxially grown quantum well structure to a quantum dot structure using the AAO
membrane as a mask.
124
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
A novel material system, InGaAsN:Sb is being explored in this framework as well for possibly
extending the GaAs based material system into the infrared optical telecommunication wavelengths.
This potentially important material only proposed in 1996 has developed to a point where good
quantum well materials can be grown. There have been a few demonstrations of its potential as a
optical source. By using the AAO templating technique as described above, explorations into
quantum dots can take place immediately and would not have to wait for developments relying on
self-assembly techniques that would take considerable amount of efforts in optimizing the growth
conditions. A study of the optical properties of InGaAsN:Sb/GaAs QDs fabricated through the
nanopore templating and etching technique has already been investigated [58]. The MBE grown
quantum well structure was first patterned with Ni nanodots to act as the mask for the subsequent RIE
in a BCl3 environment. The resultant structure is a hexagonal, periodic array of 55 nm diameter,
6.7 nm thick InGaAsN:Sb quantum dots sandwiched between GaAs layers. Unlike all other studies
normally done at low temperatures, room temperature photoluminescence was demonstrated in the
etched quantum dot array with free exposed surface and found to exceed as-grown InGaAsN:Sb/GaAs
quantum wells by several times. Stronger hole localization (in the x–y plane) in the dots appears to be
the key to the observed improvement. These findings might be particularly important for the
fabrication of the improved long wavelength room temperature operating 3D light emitting electronic
devices.
The strategy adopted in the previous references represents an effective way to fabricate highly
ordered, multi-layers of quantum dots perfectly stacked on top of each other. It also presents a simple
method to study nanopillar heterojunctions since the heterojunction interface is of high quality
proven with many years of refined epitaxial growth techniques. It is more attractive than trying to
grow good material in the nano-scale sizes and hoping for good heterojunctions. Although Mei et al.
[59] did realize MBE grown highly ordered InGaAs/GaAs and GaAs/AlGaAs quantum dot arrays
using the AAO template as a shadow mask, it is more like a point contact array on a continuum in
terms of heterojunction.
A viable approach to form QDs embedded in a wider bandgap surrounding is to first form nanocavities by the template-assisted etch followed by epitaxial regrowth in molecular-bean epitaxy
(MBE) and metalorganic chemical vapor deposition (MOCVD) processes [60,61]. It was shown that
InAs quantum dots can be selectively grown at the bottom of honeycomb hollows formed by
anodization of GaAs substrate. The average size of the InAs quantum dots and its variation were
considerably affected by the regularity of the hollows [61]. Not surprisingly, the honeycomb
hollows of anodized GaAs are of less uniformity and ordering than that of the AAO template. The
size of the hollows is of the order of 200 nm. The AAO template approach of fabricating
semiconductor anti-dot array offers a simple and effective way for optimizing the control of the size
and density of QDs on a pre-patterned substrate through the control of the diameter, spacing, and
depth of the pre-patterning.
Having semiconductor nanopillars either with homogeneous or heterogeneous material
fabricated by two distinctive methods, etching or growth, offer an opportunity to study the effects
of surface states on the properties of semiconductor nanostructures. In the conventional material
systems, bulk states dominate the overall properties. When it comes to nanostructures with confined
dimensions, such as those realized by this template method, the surface to volume ratio is so high,
surface layer properties are no longer of secondary importance. When created by dry etching, there
are plenty of dangling bonds associated with and therefore depletion layer in the surface region of
atoms. When formed by growth, the surface atoms assemble in a way to try to minimize the total
energy by forming self-closed orbitals (bonds) thereby deforming or restructuring the lattice. Thus,
with the same material of similar confined dimensions but fabricated differently, we are permitted
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
through the manipulation of surface states to create and study new properties in the nanostructures,
and in their interfaces with other materials.
4.5. Vapor–liquid–solid growth method
Porous anodic aluminum oxides are particularly useful for optical studies of the nanoparticle
structures prepared within their pores because the host oxide is transparent over much of the visible
and infrared spectral regions. Furthermore, polarization dependent optical properties of nanowires
can be determined due to the parallel orientation of the cylindrical nanopores. Ordered
semiconductor ZnO nanowire arrays have been grown in the AAO template along the nanochannels
[62]. After electrodeposition of Zn and with the metal sitting at the bottom of the AAO template, the
sample was then heated in air at 300 8C for 35 h transforming the metal into polycrystalline ZnO of
various orientations resembling standard powder diffraction patterns of ZnO. The diameters of the
nanowires ranged from 15 to 90 nm dependent on the nanopore size and the height determined by the
height of the electrodeposited Zn.
For many applications, the nanowire needs to extend beyond the template surface or be
completely removed from the AAO membrane. With the electrodeposition and oxidation method
previously described, this cannot be done. Recently, there has been more interest in a technique first
proposed in the 1960s, namely, the vapor–liquid–solid (VLS) growth method [63]. Here, metal
nanodots are used as catalysts promoting nucleation at each site forming nanowires. The VLS
reaction has been used to obtain Si [64,65], InP [66], CdS [5], and GaN [67–69] nanowires to name a
few. Extending the VSL growth method by starting with highly ordered metal nanodot arrays, a large
number of free standing highly ordered and uniform semiconductor nanowire arrays can be
synthesized with high aspect ratios covering a large area. ZnO nanowires have recently been grown
using this growth method with a high degree of ordering by using the catalytic Au nanodots as the
nucleation sites [70]. ZnO is of much interest as it is a semiconductor that has the potential for
application in a wide range of fields due to its unique properties. The material has a direct bandgap
of 3.37 eV at room temperature and its larger exciton binding energy of 60 meV and higher optical
gain, when compared to its GaN counterpart, leading to demonstrations of room temperature lasing
from single nanorods [71]. Taking advantage of its strong UV absorption, ZnO is seen as a promising
material for solar cells even in its bulk form, but with a packing density of 1011 cm2 vertically
aligned nanorods dramatically increasing the surface to volume ratio, that vision can only be
enhanced. Besides its optical applications, ZnO based surface acoustic wave devices have been
shown, exploiting its piezoelectric property. Bulk ZnO has been used as a gas sensor material relying
on the interaction of the gas molecule and the surface of the ZnO. Again, enhanced device sensitivity
can be envisioned with increased surface to volume from arrays of nanorods.
However, while many different routes to synthesize ZnO nanowires have been undertaken
including physical vapor deposition [72], vapor-phase transport [73–75], and chemical vapor
deposition [76], high fabrication costs, reduced crystallinity, wool-like geometry along with low
control over the nucleation site and size of the nanorods severely limit the development of ZnO
nanowires into useful electronic, optical, and mechanical nanodevices. Selected catalytic vaporphase transport growth methods [73], template-assisted methods [62], catalysis-driven MBE growth
methods [77], and metalorganic vapor-phase epitaxy (MOVPE) growth methods [78] have produced
site-specific nucleation of ZnO. However, for nanodevice integration, along with site-specific
nucleation is the need for the ability to place the nucleation site. By controlling the spatial
arrangement of the catalysts. Direct size and placement control of ZnO nanowires can be achieved
by using the AAO template to create an array of Au nanodots (catalysts). The diameter of the
125
126
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
nanowire is determined by the diameter of the catalyst, and thus the AAO templating method is an
ideal candidate as the diameter of the nanopores can be altered by changing electrochemical solution
used in the anodization process. By reducing the variations in the physical properties associated with
different diameters and randomness in their positioning further increasing the ease and flexibility in
engineering nanodevice characteristics.
Starting with Au nanodots 50 nm in diameter and spaced 110 nm apart center-to-center on the
GaN surface after the AAO template has been removed, the patterned GaN substrate was placed onto
a ceramic boat downstream from the ZnAs2 source supplying zinc in a horizontal tube furnace. In the
sample shown in Fig. 20, a growth time of 20 min produced nanorods of 400 nm in average length.
The average diameter of the resultant nanorods is 60 nm. The nanorods are vertically aligned with
respect to the surface of the GaN and orderly spaced in the hexagonal pattern defined the AAO
template deposition of the Au catalysts.
As shown in Fig. 20, the Au catalyst remains at the top of each ZnO nanorod indicative of the
vapor–liquid–solid growth process [63]. First, formation of an eutectic alloy droplet occurs at each
catalytic site, followed by the nucleation and growth of the solid ZnO nanorod due to the
supersaturation of the liquid droplets. Incremental growth of the nanorod takes place at the droplet
Fig. 20. ZnO nanowires grown perpendicularly to the GaN substrate: (a) top view showing hexagonal shaped nanowires
with Au catalysts remaining atop and (b) oblique angle view.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
interface constantly pushing the catalyst upwards. Thus, this growth method inherently provides sitespecific nucleation at each Au catalytic site. The presence of a metal catalyst provides a natural
contact on one end of the nanorod, which can be beneficial in certain applications where making
contacts to the nanorod is necessary.
Since GaN and ZnO share a common crystal structure with a lattice mismatch of less than
0.08%, the nanorods grown on GaN are hexagonally faced strongly suggesting that they are of
hexagonal wurzite crystal structure with peaks in the X-ray diffraction (XRD) pattern indexed to
bulk ZnO. It suggests c-axis orientation growth or perpendicular to the surface of the GaN. The band
alignment between GaN and ZnO is such that a type-II heterojunction is formed at the interface and
the associated charge transfer provides effective doping. Photoluminescence spectroscopy on ZnO
nanorods dispersed in solutions showed strong emission in the ultraviolet region with a peak at
377 nm and a FWHM of 20 nm attributed to the exciton peak. A secondary broader band in the
visible green region around 511 nm also appeared in the spectra attributed to the recombination of
photogenerated holes with the singly ionized oxygen vacancies in the ZnO crystal structure [79].
4.6. Other in-template growth methods
In recent years, nanostructured electrode materials have attracted attention since the capacity of
the electrode is critically affected by its morphology. Better rate capabilities can be obtained due to
the reduction in the distance over which Liþ ions must diffuse in the nanostructured electrodes of
rechargeable lithium-ion batteries. In addition, the surface to volume ratio of the nanopillar
electrodes is much larger than conventional unstructured electrodes leading to smaller current
densities for a given current during charging and discharging. Sol–gel chemistry is a technique that
has been employed to fabricate LiNiO2 [80], LiMn2O4 [81], LiCoO2 [82], and LiNi0.5Co0.5O2 [83]
in-template nanowires. The growth method typically involves the hydrolysis of a solution of a
precursor molecule to obtain a suspension of colloidal particles. This is the sol part of the equation.
The gel is composed of aggregated sol particles thermally treated to yield the desired material. To
obtain in-template nanowires, the AAO template is dipped into the sol allowing the solution to
diffuse into the nanochannels. After a set length of time, the template is removed from the sol and
heated at elevated temperatures. Fig. 21 shows SEM images of LiMn2O4 nanowires. Highly ordered
zirconia nanowire arrays have also been demonstrated by the AAO template method using sol–gels
[84].
Cheng et al. synthesized in-template GaN nanowires through a gas-phase reaction of Ga2O
vapor with flowing NH3. While Yu et al. fabricated in-template poly(p-phenylene) PPP nanowire
arrays in also a gas-phase reaction of benzene, aluminum chloride anhydrous, and cupric chloride
anhydrous [85]. PPP is an interesting material because of its capability of being transformed from an
electrical insulator to an electrical conductor with applications in light-weight rechargeable batteries
and blue emitting devices.
Clearly from the numerous examples, the AAO templating method can be seen as an extremely
versatile tool in the fabrication of highly ordered, densely packed nanowires and nanopillars of
various materials.
5. Nanopores and anti-dot arrays
The versatility of the non-lithographic AAO template approach allows it to be extended into
fabricating superlattice nanopores. Much of the attention in research has gone into nanowires,
127
128
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 21. LiMn2O4 nanowires fabricated using sol–gel techniques from [81]: (a) magnified cross-section view and (b) crosssection view showing length of nanowire about 30 mm.
nanopillars, and nanodots. This natural course of progression from bulk materials to nanowires to
nanodots in an attempt to exploit the quantum properties of small sizes is indeed fascinating and well
deserving of the attention, however, work with ‘‘empty’’ spaces is proving to be just as interesting as
new physics and phenomenons are being uncovered. In the sections to follow, some of the novel
physics utilizing nanopores will be presented.
The fabrication of nanopores is a relatively simple process. In a ‘‘positive’’ transfer, the
nanopore array pattern is replicated conveniently through etching, chemically and/or physically in
the substrate resulting in lateral superlattices of what are often referred to as ‘‘anti-dots’’ (pores).
One procedure we often follow employs the AAO nanopore array membrane as an etch mask.
Briefly, it goes like this: After the template has been formed and the barrier layer removed as
described in previous sections, the highly ordered AAO nanopore film is placed onto the desired
substrate. Dry etch methods, such as plasma etching, ion milling and reactive ion etching, are
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 22. GaAs nanopore array from [86]: (a) cross-sectional view and (b) plot shows the 10:1 aspect ratio of the nanopore
depth to diameter greatly reduces the etch rates inside the nanopores.
used for transferring the pattern to the substrate because of their better directionality over wet
etching.
Nakao et al. were amongst the first to report the formation of nanopore arrays in semiconductors
using the AAO templating method in 1998 [86]. GaAs and InP nanopore arrays were formed using a
reactive beam etching technique with Br2–N2 gases at elevated substrate temperatures. Nanopores of
up to 1 mm in depth were etched, however, as noted by the authors, the holes at this depth have partly
collapsed. Fig. 22(a) shows the cross-section SEM image of a 80 nm diameter and a 500 nm deep
GaAs nanopore array. At these high aspect ratio approaching 10:1 (depth to width), etching rates
inside the pores are dramatically reduced to the tune of one order of magnitude slower as shown in
the graph in Fig. 22(b). Typically thickness used are around 500 nm still leaving an aspect ratio of
129
130
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
10:1 with respect to the width of each nanochannel. When one takes in account the 0.5 mm AAO
template which is being etched, albeit at a much slower etch rate, the aspect ratio at the end of
nanopore etching process is almost 20:1. Thus, an argument is made for reducing the thickness of the
AAO template to as thin as possible if deep nanopores are desired with it still being possible to
physically handle. Micro-loading effects were also seen as 60 nm diameter pores were 10%
shallower in depth than 80 nm diameter pores.
The atomic force microscope (AFM) images in Fig. 23 show a nanopore array on GaN with a
depth of about 120 nm, a mean pore diameter of approximately 40 nm and a period of 100 nm by
Fig. 23. AFM images of a nanopore array in GaN.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
RIE. The typical nanopore etching conditions for different semiconductors such as silicon [87] are
the same as for etching nanopillars listed in Table 2. Guo et al. also employed reactive ion etching
techniques using a combination of CH4 and H2 gases to fabricate 60 nm nanohole channels in ZnTe
[88].
In transferring the AAO nanopore pattern onto the substrate, reactive ion etching is typically
used due to its directive etching nature, its smooth etched surface morphology, high etch rates, and
vertical sidewalls. In this process, any gap between the template and the substrate should be
minimized to ideally a direct contact so as to prevent unwanted etch along the gap. To circumvent the
challenges in placing the ultra-thin AAO film onto the receiving substrate, an Al film can be
sputtered or evaporated onto the substrate as done by Crouse et al. [89]. After evaporating 2 mm of Al
onto a conducting Si substrate, anodization using similar recipes proceeded. One limitation to this
process is the substrate must be conducting as the anodization procedure requires the voltage drop on
the Al surface. After anodization, the AAO film was thinned from 2 to 300 nm by a 1 h argon-ion
milling procedure to make the subsequent RIE pattern transfer to Si substrate easier. Fig. 24(a) shows
the AAO layer on the Si substrate after anodization, and Fig. 24(b) the result of a 1 min RIE into Si
with the ordered AAO nanopore film still attached to the substrate.
Fig. 24. Anodization performed with Al film on Si from [89]: (a) cross-sectional view of highly ordered AAO
nanochannels and (b) after nanopore pattern transfer.
131
132
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
5.1. Optical properties and phenomenons
Arrays of nanopores with a high degree of uniformity and ordering have lead to new effects just
as important and interesting in the collective excitation or relaxation modes and in the near-field (or
sub-wavelength) coupling regime. The systems in such collective modes can manifest behaviors
deviating fundamentally from the classic processes such as diffraction limits and photoluminescence
[90]. Ebbesen et al. demonstrated that the zero-order transmission of light (the incident and detected
light are collinear) through sub-wavelength hole arrays made in a metal film can be orders of
magnitude larger than is expected from classical aperture theory [91]. Sharp peaks are observed in
the transmission spectra at wavelengths as large as ten times the diameter of the cylindrical
nanopores. From classical theory, the transmission of a sub-wavelength aperture is proportional to
the fourth power of the ratio of its diameter and wavelength of light. Thus, one would expect
transmission in the order of 1% efficiency instead of the observed ‘‘apparent’’ 130% when
normalized to the area covered by holes [92]. This extraordinary phenomenon can be used in a
number of applications such as sub-wavelength photolithography, near-field microscopy, molecular
imaging, and data storage. However, the underlying physical mechanism has not yet been fully
understood despite some elucidations based on numerical models [93]. The few efforts to date all
point to the central role played by surface plasmons and their collective excitations by and
interactions with photons. Although the experiments have been fewer than desired, they are very
exciting for they represent a new opportunity in realizing the long-sought coherent and collective
interactions of electrons with a radiation field. Progress to date on this new frontier has been
confined to samples made by e-beam lithography or focused ion-beam milling, but more rapid
progress can certainly be made with strategies readily implementable with the AAO nanopore array
such as simply coating the membrane by evaporating metal film of various thickness and materials.
Although the use of AAO templates have not been demonstrated for this application, it is not difficult
to envision its place in this field of work. Recently, metal hole arrays were fabricated by a two-step
molding process using the AAO template as the medium for the high aspect ratio holes. Instead of
relying on the two-step anodization process to form the ordered nanopore pattern, an e-beam written
SiC mold was used to imprint the Al replacing the first anodization step [94]. With the Al
pretextured with shallow concave imprints, the standard anodization procedure followed creating
highly ordered nanopores with an aspect ratio of 7. A Ni metal film was then evaporated onto the
nanopore surface with the results shown in Fig. 25. The depth of the nanopores is 2.5 mm with a pore
diameter of 350 nm and a pore spacing of 500 nm.
Nanopatterned surfaces could exhibit very low reflectivity over a broad bandwidth and a very
large viewing field, if the nanopores or nanopillars are etched deep enough into the surface.
Conceptually, the large aspect ratio and sub-wavelength texture create an effective gradual index
change from that of air at the top to that of the substrate at the bottom. The weak but numerous
distributed scatterings by the sub-wavelength features also gives rise to diffusive optical transport,
adding to the effective grading of the index. Theoretical and experimental evidence supporting this
expectation were first obtained recently in silicon where a reflectivity as low as 1.6% was measured
over a range of 300–800 nm [95]. The reflectivity of a polished silicon wafer over the same
wavelength range is about 35%. The non-lithographic template fabrication approach is uniquely
advantageous for patterning curved surfaces for experimental demonstration and optimization of the
expected broadband anti-reflection characteristics, and for enabling novel coherent electron transport
studies. This method is also readily scalable to large areas.
Silicon-based photonics has been much sought after with great anticipation of enabling direct
integration with its fully developed electronic counterpart. Silicon photodetection can be readily
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Fig. 25. Ni film with highly ordered nanopores produced from e-beam written SiC mold, imprinting, followed by
anodization: (a) top view and (b) cross-sectional view from [94].
performed, however, photon emission from bulk silicon is approximately 1% efficient at best due to
the fact it is an indirect bandgap semiconductor. By terminating the continuum of the silicon crystal
with interfaces nanometers apart, one can both break the crystal symmetry and introduce the
quantum confinement, and thereby, with the former, alter the dispersion relation, and with the latter,
enhance the exciton and optical oscillator strength. A quantum anti-dot configuration, where periodic
nanopores are created in a thin layer of silicon, is electrically connected with constrictions that are
narrow (quantum) enabling possible cascade emission and amplification. The anti-dot configuration
also helps with the desire of minimizing re-absorption enabling the possibility of a silicon laser.
On another front where the creation of nanopore array on semiconductor is followed by
subsequent epitaxial growth, positive but still preliminary results have been obtained. Exploratory
studies of nano-heteroepitaxial growth of GaN films on highly ordered nanoporous Si(1 1 1)
substrates show that a multi-fold increase in photoluminescence intensity can be obtained. The
observed improved quality suggests that the uniformly nanopatterned interface structure, as
expected, does offer an advantage for releasing stress due to lattice and thermal expansion
differences. The nanopore array surface with the pore diameter of 60 nm and periodicity of 110 nm
exhibits significant benefits for emissivity, film growth, and the relaxation of misfit and thermal
strain in the film. Peak shifts in photoluminescence and Raman spectroscopy indicate that the
material grown on nanopores is more relaxed than films grown on flat Si(1 1 1) substrates. The
effects of nanopore topography on the nucleation of GaN films offer a potential path to significant
improvement of III-nitride heteroepitaxy for device applications [96].
133
134
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
5.2. Superconducting nanomesh
This simple approach of using the highly ordered AAO template to fabricating nanomeshes has
already led to significant advances in vortex physics of normal superconductors [97]. In a recent
experiment, AAO membrane was covered with a 100 nm film of Nb deposited by magnetron
sputtering followed by a 10 nm thick Ag capping layer as shown in Fig. 26(a). The SEM image
shows the highly ordered nanopore pattern in the cleaved AAO membrane after Nb and Ag layers
have been deposited. Since the sputtered Nb and Ag layers deposit onto the walls of the AAO
nanopores, the pore opening diameter is reduced to 45 nm. The pore spacing is left unchanged
at 101 nm. The pore density in these arrays is almost two orders of magnitude greater than in
typical lithographically prepared micron-sized hole arrays. Magnetization and magneto-transport
Fig. 26. (a) Oblique angle view of an AAO membrane covered with 100 nm of Nb capped with 10 nm of Ag. The white
line shows a grain boundary between hexagonal nanopore patterns of different orientation and (b) magnetic moment of the
Nb-AAO array for a perpendicular applied field. The matching fields are denoted H1, H2, and H3.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
measurements reveal effects not seen in the lithographically produced hole arrays as the higher
density non-lithographically produced hole arrays puts the templated vortex array into the strongly
interacting limit where vortex–vortex interactions reinforce the natural pinning of the template.
Revealed effects include:
Confinement effects where the dimensional crossover from 2D superconductivity behavior to 1D
behavior are seen above 6 K.
Little-Parks oscillations of the upper critical field are seen above 6 K.
The presence of strong pinning at low temperatures. The matching fields and strong pinning reveal
pronounced stepwise variation of the critical current at multiples of the fundamental matching
field and extends to much higher fields, 7.5 kG for the third matching field, whereas the
matching field effects decay above a few hundred Oe; see Fig. 26(b).
Pinning is enhanced by two orders of magnitude over that in continuous Nb films. The strong
pinning enables even low pinning superconductors to carry a high critical current
Viewing this work from a broader perspective, one can see that similar strong collective
behaviors arising from coherent couplings of electronic and optical modes induced and modulated
by periodic distributed near-field coupling can also be expected in non-superconducting media.
6. Summary and future prospects
We have attempted in this review to summarize the progress of a non-lithographic technique
using the self-organized, highly ordered AAO membrane as a template to fabricate various
nanometric structures and briefly present their immediate contributions to nanotechnology in a
field still in its infancy. The AAO membrane itself determines the uniformity, packing density, and
size of the nanometric superlattices made by using this approach. By varying the anodization
conditions, the diameter of the nanopores of the AAO templates can be altered between 20 and
120 nm. The packing density can reach the order of 1012/cm2 covering an area that is readily
scalable. The fact that material systems as different as metals, semiconductors, and carbon
nanotubes can be treated with the same ease and the fact that structural forms as different as dots,
tubes, pillars, anti-dots (nanopore meshes) can be fabricated with the same base technology attests
to the generality of this nano-fabrication approach. The flexibility of using a variety of materials,
the accurate control over fabrication process, and the command over nanopore template features
give us the freedom of engineering various physical properties determined by the shape, size,
composition, and doping of the nanostructures as well as many new properties enabled by their
couplings. The novel nanometric superlattice platform realized by this unique nano-fabrication
approach can powerfully enable a broad range of applications in electronics, optics, mechanics, and
bio-technologies. Some of these are extraordinary properties and functions not intrinsic to the
individual nano-elements but only arise from the collective behaviors of the large ensemble of
periodically coupled and uniform nanostructures due to the near-field coupling and sub-wavelength
spacing.
By manipulating the geometry of otherwise conventional materials into within the mesoscopic
length scales comparable to the mean free lengths, exciton radius, or Fermi wavelength, we can
realize unconventional physical properties due to the size effects as well as due to surface
reconstructions. When combined with compositional changes on the same length scales, nano-scale
engineering offers us a truly new capability in science and technology.
135
136
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
Acknowledgements
This paper covers a large set of results and findings from a number of projects supported by
AFOSR, AF Labs, CIAR, DARPA, Motorola, NSF, Nortel, and ONR. Some of the specifics included
in this review were summarized from the works reported in various publications and presentations
(see references) by our colleagues, including Jianyu Liang, Chris Papadopoulos, Aijun Yin, Baohe
Chang, Nikolai Kouklin, and Jing Li.
References
[1]
[2]
[3]
[4]
[5]
[6]
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
[15]
[16]
[17]
[18]
[19]
[20]
[21]
[22]
[23]
[24]
[25]
[26]
[27]
[28]
[29]
[30]
[31]
[32]
[33]
[34]
[35]
[36]
[37]
[38]
S.J. Tan, A.R.M. Verschueren, C. Dekker, Nature 393 (1998) 49.
H.W.C. Postma, T. Teepen, Z. Yao, M. Grifoni, C. Dekker, Science 293 (2001) 76.
Y. Huang, X. Duan, Y. Cui, L.J. Lauhon, K.H. Kim, C.M. Lieber, Science 294 (2001) 1313.
A. Bachtold, P. Hadley, T. Nakanishi, C. Dekker, Science 294 (2001) 1317.
X. Duan, Y. Huang, R. Agarwal, C.M. Lieber, Nature 421 (2003) 241.
F. Keller, M.S. Hunter, D.L. Robinson, J. Electrochem. Soc. 100 (1953) 411.
H. Masuda, H. Yamada, M. Satoh, H. Asoh, M. Nakao, T. Tamamura, Appl. Phys. Lett. 71 (1997) 2770.
H. Masuda, F. Hasegwa, S. Ono, J. Electrochem. Soc. 144 (1997) L127.
O. Jessensky, F. Muller, U. Gosele, Appl. Phys. Lett. 72 (1998) 1173.
A. Despic, V.P. Parkutik, in: J.O. Bockris, R.E. White, B.E. Conway (Eds.), Modern Aspects of Electrochemistry,
Plenum Press, New York, 1989, p. 401.
J.P. O’Sullivan, G.C. Wood, J. Electrochem. Soc. 116 (1969) 1351.
M. Konno, M. Shindo, S. Sugawara, S. Saito, J. Membr. Sci. 37 (1988) 193.
H. Masuda, K. Fukuda, Science 268 (1995) 1466.
S. Bandyopadhyay, A.E. Miller, H.C. Chang, G. Banerjee, V. Yuzhakov, D.F. Yue, R.E. Ricker, S. Jones, J.A.
Eastman, E. Baugher, M. Chandrasekhar, Nanotechnology 7 (1996) 360.
O. Jessensky, F. Muller, U. Gosele, J. Electrochem. Soc. 145 (1998) 3735.
A.P. Li, F. Muller, A. Birner, K. Nielsch, U. Gosele, J. Appl. Phys. 84 (1998) 6023.
T. Gao, J.C. Fan, G.W. Meng, Z.Q. Chu, L.D. Zhang, Thin Solid Films 401 (2001) 102.
H. Masuda, M. Sotoh, Jpn. J. Appl. Phys., Part 2 35 (1996) L126.
J. Liang, H. Chik, A. Yin, J.M. Xu, J. Appl. Phys. 91 (2002) 2544.
X. Mei, M. Blumin, D. Kim, Z. Wu, H.E. Ruda, J. Cryst. Growth 251 (2003) 253.
H. Masuda, K. Yasui, K. Nishio, Adv. Mater. 12 (2000) 1031.
R.W. Keyes, Science 230 (1985) 138.
C.P. Collier, E.W. Wong, M. Belohradsky, F.M. Raymo, J.F. Stoddart, P.J. Kuekes, R.S. Williams, J.R. Heath,
Science 285 (1999) 391.
R.P. Cowburn, M.E. Welland, Science 287 (2000) 1466.
G. Sauer, G. Brehm, S. Schneider, K. Nielsch, R.B. Wehrspohn, J. Choi, H. Hofmeister, U. Gosele, J. Appl. Phys. 91
(2002) 3243.
S. Yang, H. Zhu, D. Yu, Z. Jin, S. Tang, Y. Du, J. Magnestism Magnetic Mater. 222 (2000) 97.
D. Routkevitch, A.A. Tager, J. Haruyama, D. Almawlawi, M. Moskovits, J.M. Xu, IEEE Trans. Electron Devices 43
(1996) 1646.
M. Zheng, L. Menon, H. zeng, Y. Liu, S. Bandyopadhyay, R.D. Kirby, D.J. Sellmyer, Phys. Rev. B 62 (2000)
12283.
X.F. Wang, J. Zhang, H.Z. Shi, Y.W. Wang, G.W. Meng, X.S. Peng, L.D. Zhang, J. Fang, J. Appl. Phys. 89 (2001)
3847.
D. Routkevitch, T. Bigioni, M. Moskovits, J.M. Xu, J. Phys. Chem. 100 (1996) 14037.
J. Haruyama, D.N. Davydov, D. Routkevitch, D. Ellis, B.W. Statt, M. Moskovits, J.M. Xu, Solid-State Electron. 42
(1998) 1257.
D. Xu, Y. Xu, D. Chen, G. Guo, L. Gui, Y. Tang, Chem. Phys. Lett. 325 (2000) 340.
A.J. Yin, J. Li, W. Jian, A.J. Bennett, J.M. Xu, Appl. Phys. Lett. 79 (2001) 1039.
S.Y. Chou, M.S. Wei, P.R. Krauss, P.B. Fischer, J. Appl. Phys. 76 (1994) 6673.
K. Nieisch, R.B. Wehrspohn, J. Barthel, J. Kirschner, U. Gosele, S.F. Fischer, H. Kronmuller, Appl. Phys. Lett. 79
(2001) 1360.
S. Iijima, Nature 354 (1991) 56.
C.T. White, D.H. Robertson, J.W. Mintmire, Phys. Rev. B 47 (1993) 5485.
J.W.G. Wildoer, L.C. Venema, A. Rinzler, R.E. Smalley, C. Dekker, Nature 391 (1998) 59.
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
[39]
[40]
[41]
[42]
[43]
[44]
[45]
[46]
[47]
[48]
[49]
[50]
[51]
[52]
[53]
[54]
[55]
[56]
[57]
[58]
[59]
[60]
[61]
[62]
[63]
[64]
[65]
[66]
[67]
[68]
[69]
[70]
[71]
[72]
[73]
[74]
[75]
[76]
[77]
[78]
[79]
[80]
[81]
[82]
[83]
[84]
[85]
[86]
[87]
[88]
[89]
[90]
[91]
[92]
R.V. Parthasarathy, K.L.N. Phani, C.R. Martin, Adv. Mater. 7 (1995) 896.
J. Li, T. Haslett, M. Moskovits, Chem. Mater. 10 (1998) 1963.
T. Koytani, L. Tsai, A. Tomita, Chem. Mater. 8 (1996) 2190.
W.Z. Li, S.S. Xie, L.X. Qian, B.H. Chang, B.S. Zou, W.Y. Zhuo, R.A. Zhao, G. Wang, Science 274 (1996) 1701.
Z.F. Ren, Z.P. Huang, J.W. Xu, J.H. Wang, P. Bush, M.P. Siegal, P.N. Provencio, Science 282 (1998) 1105.
T. Kyotani, L. Tsai, A. Tomita, Chem. Mater. 8 (1996) 2109.
J. Li, C. Papadopoulos, J.M. Xu, Appl. Phys. Lett. 75 (1999) 367.
W. Hu, D. Gong, Z. Chen, L. Yuan, K. Saito, C.A. Grimes, P. Kichambare, Appl. Phys. Lett. 79 (2001) 3083.
Y.C. Sui, D.R. Acosta, J.A. Gonzalez-Leon, A. Bermudez, J. Feuchtwanger, B.Z. Cui, J.O. Flores, J.M. Saniger, J.
Phys. Chem. B 105 (2001) 1523.
C. Papadopoulos, B.H. Chang, A.J. Yin, J.M. Xu, Inter. J. Nanoscience 1 (2002) 205.
W.B. Choi, J.U. Chu, K.S. Jeong, E.J. Bae, J.W. Lee, J.J. Kim, J.O. Lee, Appl. Phys. Lett. 79 (2001) 3696.
P.G. Collins, H. Bando, A. Zettl, Nanotechnology 9 (1998) 153.
J. Li, C. Papadopoulos, J.M. Xu, Nature 402 (1999) 253.
C. Papadopoulos, A. Rakitin, J. Li, A.S. Vedeenev, J.M. Xu, Phys. Rev. Lett. 85 (2000) 3476.
J. Liang, H. Chik, J.M. Xu, IEEE J. Select. Top. Qunatum Electron. 8 (2002) 998.
K. Mukai, Y. Nakata, K. Otsubo, M. Sugawara, N. Yokoyama, H. Ishikawa, IEEE J. Quantum Electron. 36 (2000)
472.
Q. Xie, J.L. Brown, R.L. Jones, J.E. Van Nostrand, K.D. Leedy, Appl. Phys. Lett. 76 (2000) 3082.
N.N. Ledentsov, M. Grundmann, F. Heinrichsdorff, D. Bimberg, V.M. Ustinov, A.E. Zhukov, M.V. Maximov, Z.I.
Alferov, J.A. Lott, IEEE J. Select. Top. Qunatum Electron. 6 (2000) 439.
L. Chen, A. Yin, J.X. Im, A.V. Nurmikko, J.M. Xu, J. Han, Phys. Stat. Sol. A 188 (2001) 135.
N. Kouklin, H. Chik, J. Liang, M. Tzolov, J.M. Xu, J.B. Heroux, W.I. Wang, J. Phys. D: Appl. Phys. 36 (2003)
2634.
X. Mei, D. Kim, H.E. Ruda, Q.X. Guo, Appl. Phys. Lett. 81 (2002) 361.
M. Lachab, M. Nozaki, J. Wang, Y. Ishikawa, Q. Fareed, T. Wang, T. Nishikawa, K. Nishino, S. Sakai, J. Appl. Phys.
87 (2000) 1374.
Y. Morishita, J. Sunagawa, Y. Yumoto, S. Kawai, J. Cryst. Growth 227 (2001) 1049.
Y. Li, G.W. Meng, L.D. Zhang, F. Phillipp, Appl. Phys. Lett. 76 (2000) 2011.
R.S. Wagner, W.C. Ellis, Appl. Phys. Lett. 4 (1964) 89.
Y. Wu, P. Yang, J. Am. Chem. Soc. 123 (2001) 3165.
J. Westwater, D.P. Gosain, S. Tomiya, S. Usui, H. Ruda, J. Vac. Sci. Technol. B 15 (1997) 554.
X. Duan, Y. Huang, J. Wang, Y. Cui, C.M. Lieber, Nature 409 (2001) 66.
X. Duan, C.M. Lieber, J. Am. Chem. Soc. 122 (2000) 188.
J. Zhang, L.D. Zhang, X.F. Wang, C.H. Liang, X.S. Peng, Y.W. Wang, J. Chem. Phys. 115 (2001) 5714.
J. Zhang, L. Zhang, J. Phys. D: Appl. Phys. 35 (2002) 1481.
H. Chik, J. Liang, S.G. Cloutier, N. Kouklin, J.M. Xu, submitted for publication.
M.H. Huang, S. Mao, H. Feick, H. Yan, Y. Wu, H. Kind, E. Weber, R. Russo, P. Yang, Science 292 (2001) 1897.
Y.C. Kong, D.P. Yu, B. Zhang, W. Fang, S.Q. Feng, Appl. Phys. Lett. 78 (2001) 407.
M.H. Huang, Y. Wu, H. Feick, N. Tran, E. Weber, P. Yang, Adv. Mater. 13 (2001) 113.
Y.W. Wang, L.D. Zhang, G.Z. Wang, X.S. Peng, Z.Q. Chu, C.H. Liang, J. Cryst. Growth 234 (2002) 171.
J.Y. Li, X.L. Chen, H. Li, M. He, Z.Y. Qiao, J. Cryst. Growth 233 (2001) 5.
J.J. Wu, S.C. Liu, J. Phys. Chem. B 106 (2002) 9546.
Y.W. Heo, V. Varadarajan, M. Kaufman, K. Kim, D.P. Norton, F. Ren, P.H. Fleming, Appl. Phys. Lett. 81 (2002)
3046.
W.I. Park, D.H. Kim, S.W. Jung, G.C. Yi, Appl. Phys. Lett. 80 (2002) 4232.
K. Vanheusden, W.L. Warren, C.H. Seager, D.K. Tallant, J.A. Voigt, B.E. Gnade, J. Appl. Phys. 79 (1996) 7983.
Y. Zhou, J. Huang, C. Shen, H. Li, Mater. Sci. Eng. A335 (2002) 260.
Y. Zhou, C. Shen, J. Huang, H. Li, Mater. Sci. Eng. B95 (2002) 77.
Y. Zhou, C. Shen, H. Li, Solid State Ionics 146 (2002) 81.
Y. Zhou, H. Li, J. Mater. Sci. 37 (2002) 5261.
H. Xu, D.H. Qin, Z. Yang, H.L. Li, Mater. Chem. Phys. 80 (2003) 524.
B. Yu, H. Li, Mater. Sci. Eng. A325 (2002) 215.
M. Nakao, S. Oku, T. Tamamura, K. Yasui, H. Masuda, Jpn. J. Appl. Phys., Part 1 38 (1999) 1052.
S. Shingubara, O. Okino, Y. Murakami, H. Sakaue, T. Takahagi, J. Vac. Sci. Technol. B 19 (2001) 1901.
Q. Guo, T. Tanaka, M. Nishio, H. Ogawa, X. Mei, H. Ruda, Jpn. J. Appl. Phys., Part 2 40 (2002) L118.
D. Crouse, Y.H. Lo, A.E. Miller, M. Crouse, Appl. Phys. Lett. 76 (2000) 49.
R.H. Dicke, Phys. Rev. 93 (1954) 99.
T.W. Ebbesen, H.J. Lezec, H.F. Ghaemi, T. Thio, P.A. Wolff, Nature 391 (1998) 667.
L. Martin-Moreno, F.J. Garcia-Bidal, H.J. Lezec, K.M. Pellerin, T. Thio, J.B. Pendry, T.W. Ebbesen, Phys. Rev. Lett.
86 (2001) 1114.
137
138
H. Chik, J.M. Xu / Materials Science and Engineering R 43 (2004) 103–138
[93]
[94]
[95]
[96]
[97]
L. Salomon, F. Grillot, A.V. Zayats, F. de Fornel, Phys. Rev. Lett. 86 (2001) 1110.
K. Nishio, M. Nakao, A. Yokko, H. Masuda, Jpn. J. Appl. Phys., Part 2 42 (2003) L83.
Y. Kanamori, K. Hane, H. Sai, H. Yugami, Appl. Phys. Lett. 78 (2001) 142.
J. Liang, S.K. Hong, N. Kouklin, R. Beresford, J.M. Xu, Appl. Phys. Lett. 83 (2003) 1752.
G.W. Crabtree, U. Welp, Z.L. Xiao, J.S. Jiang, V.K. Vlasko-Vlasov, S.D. Bader, J. Liang, H. Chik, J.M. Xu, Physica
C 387 (2003) 49.