Synthesis and optical properties of II–VI 1D...

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Synthesis and optical properties of II–VI 1D nanostructures†

Muhammad Iqbal Bakti Utama,a Jun Zhang,a Rui Chen,a Xinlong Xu,b Dehui Li,a Handong Suna

and Qihua Xiong*ac

Received 29th October 2011, Accepted 18th November 2011

DOI: 10.1039/c1nr11612f

1D nanostructures from II–VI semiconductors have been demonstrated to exhibit outstanding optical

properties with strong promise for novel optoelectronic devices with augmented performance and

functionalities. Herein, we present a comprehensive review discussing important topics pertinent to the

fundamental properties and applications of II–VI 1D nanostructures. With practical applications in

mind, the considerations, principles and experimental techniques on the sample preparation of high

quality 1D nanostructures are highlighted. Fundamentals on the optical properties of II–VI materials,

along with relevant investigation techniques and recent progress in the field, are also extensively

discussed. With the steady development of their synthesis, characterization and device fabrication, it is

strongly expected that II–VI 1D nanostructures will assume a unique position in future technology.

aDivision of Physics and Applied Physics, School of Physical andMathematical Sciences, Nanyang Technological University, Singapore637371. E-mail: [email protected] Key Laboratory of Photoelectric Technology and FunctionalMaterials (Culture Base), National Photoelectric Technology andFunctional Materials and Application of Science and TechnologyInternational Cooperation Base, Northwest University, Xi’an, 710069,People’s Republic of ChinacDivision of Microelectronics, School of Electrical and ElectronicEngineering, Nanyang Technological University, Singapore 639798

† This article was submitted as part of a collection highlighting papers onthe ‘Recent Advances in Semiconductor Nanowires Research’ fromICMAT 2011.

Handong Sun

Dr Handong Sun obtained his

PhD in Physics from Hong

Kong University of Science and

Technology. He is now working

as an assistant professor in the

School of Physical and Mathe-

matical Sciences, Nanyang

Technological University

(NTU), Singapore. Before he

moved to NTU, he had been

working in the Institute of

Physical and Chemical Research

(RIKEN), Japan, and the

Institute of Photonics, The

University of Strathclyde, UK.

His current research interest is

focused on spectroscopic characterization of nanostructured

materials (inorganic, organic and hybrid structures), optoelec-

tronic devices, and plasmonics-based physics and devices.

1422 | Nanoscale, 2012, 4, 1422–1435

1. Introduction

Nanostructures, generally defined as systems in possession of at

least one dimension with a size between 1 to 100 nm, have

garnered immense attention from both curiosity-driven and

technologically-motivated research.1–6 Numerous fascinating

and groundbreaking advances have been demonstrated on the

potential applications of nanostructures, in part due to the

unique platform they provide to probe and manifest various size-

and structure-dependent material properties where the relevance

of quantum confinement effects is greater.

Qihua Xiong

Dr Qihua Xiong received his BS

degree from Wuhan University

and his MS degree from

Shanghai Institute of Applied

Physics, Chinese Academy of

Sciences. He received a PhD

degree in materials science

under the supervision of Prof.

Peter Eklund at Pennsylvania

State University in 2006. After

postdoctoral experience in Prof.

Charles Lieber’s group at Har-

vard University, he joined

Nanyang Technological Univer-

sity as a Nanyang assistant

professor in 2009. He is a Fellow

of the Singapore National Research Foundation. Prof. Xiong’s

research focuses on a range of topics in emerging nanomaterials

and nanostructures, from fundamental properties and device

applications in energy and biosensing.

This journal is ª The Royal Society of Chemistry 2012

http://dx.doi.org/10.1039/c1nr11612f






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In particular, nanostructures with quasi-one-dimensionality

(1D) have been proposed as active components in devices due to

their distinguished performance in electronics, photonics, and

optoelectronics. Among the examples of 1D nanostructures are

carbon nanotubes and semiconductor nanowires—which were

originally called ‘‘whiskers’’—along with the more recent vari-

ants such as nanobelts. However, the lack of reliable techniques

to produce uniform semiconducting carbon nanotubes has

impeded the industrial adoption of carbon nanotube technolo-

gies.2 Meanwhile, techniques to prepare high quality semi-

conducting nanowires with good control and reproducibility

have been made available by meticulous studies in the field within

the last decade.6 Exciting progress within the last few years in the

field of nanowires, as an important research front in nanotech-

nology, has strengthened the expectation of the pivotal role that

nanowires may possibly hold in the future.

II–VI semiconductors (zinc chalcogenides: ZnO, ZnS, ZnSe,

ZnTe; cadmium chalcogenides: CdS, CdSe, CdTe) are among

a class of materials that has been intensively studied in the field of

1D nanostructures.7–13 II–VI compounds are recognized as

promising materials in optoelectronics owing to their high

sensitivity and quantum efficiency.7 Additionally, the range of

their wide direct band gap implies a flexibility to fabricate

optoelectronic devices to work on various wavelengths of choice.

Nevertheless, the applications of bulk and films of II–VI semi-

conductors in industry have been hindered by the strict

requirement of epitaxial growth and substrate quality, control of

defects, and difficulty in doping as opposed to that of III–V

semiconductors. This situation is in direct contrast to 1D nano-

structures with good crystallinity and controlled doping which is

easier to prepare, facilitating fundamental studies as well as

possible device applications of II–VI nanostructures. This review

will therefore be focused on the preparation techniques of II–VI

1D nanostructures, the fundamental of optical properties rele-

vant to II–VI materials, and their applications in photonic and

optoelectronic devices.

2. Synthesis

The preparation of 1D nanostructures from II–VI semi-

conductors is a crucial stage for both the study of fundamental

properties and applications of the materials toward novel tech-

nologies. Thus, as a prerequisite, it is important to develop and

utilize reliable methods with the ability to produce well

controlled dimensions, morphology, and phase purity of nano-

structures with the desired functionality. In this regard, synthesis

approaches have been thoroughly explored and proposed as

a potential alternative to the production of nanostructures from

the more mature top-down fabrication techniques of the semi-

conductor industry. The advanced lithographic techniques of

top-down fabrication have not been able to produce nano-

materials with high throughput, in addition to the escalating

production and equipment costs for further structural minimi-

zation. Meanwhile, various synthesis techniques, adopting

a bottom-up paradigm by constructing structures from smaller

building blocks, have been reported to be able to generate 1D

nanostructures with high throughput at reasonable cost effec-

tiveness and good reproducibility.2,5 Nevertheless, the capability

of top-down processes is not to be depreciated, as the integration


of nanostructures into functional and scalable devices may still

necessitate the strong organization and spatial control of

lithography.

Synthesis of 1D nanostructures is performed by promoting

nucleation, crystallization, and growth in only one direction

while suppressing the 2D growth. The various strategies that

have been demonstrated for the growth of II–VI 1D nano-

structures can be categorized according to the reaction media

and precursors that are used during the synthesis process, which

is either a solution or vapor based technique.3 The great benefits

of vapor based techniques lie in their ability to produce very high

quality materials due to the high temperature involved in the

growth process, which allow the building blocks of the growth to

rearrange themselves to develop a long-range ordered crystalline

lattice.5,7 Furthermore, the experimental methods and technol-

ogies of vapor based techniques are already established, due to

their intensive use in semiconductor research and in industry,

and thus are compatible with device integration. In comparison,

solution based techniques allow a relatively lower temperature

(T < 400 �C) to produce nanostructures and is auspicious from

the economic standpoint.14,15 However, a major demerit of most

of the solution approaches is their difficulty to synthesize well-

crystallized nanowires by achieving simultaneously high

throughput, well-defined morphology and orientation, and

positioning control, thus strongly compromising their integra-

bility into optoelectronic devices and/or productivity.

In the following, we will discuss briefly the growth mechanism

and basic principles of the experimental techniques that have

been demonstrated to be generally applicable in producing high

quality II–VI 1D nanostructures while referring to representative

examples, whenever available.

2.1. Growth mechanism

The growth of II–VI 1D nanostructures by the vapor phase

typically follows the combination of two possible underlying

growth mechanisms: the vapor–liquid–solid and vapor–solid

mechanisms. A unique case of self-catalytic growth has also

received much attention which warrants further discussion in this

review.

2.1.1. Vapor–liquid–solid (VLS) growth. VLS growth16 is

well-accepted and among the most widely attributed mechanisms

for the growth of nanowires. An in situ transmission electron

microscopy (TEM) observation of nanowire growth with the

mechanism was demonstrated, which allowed an understanding

of the growth stages.17 The kinetics of the growth have also been

extensively examined.18–21 A schematic of VLS growth is shown

in Fig. 1. VLS growth is promoted by the presence of a metallic

particle droplet, whose role resembles the function of a catalyst in

a reaction, to direct the 1D growth. Thus, VLS growth is often

referred to as ‘‘catalytic’’ growth. The metallic particle then forms

a liquid alloy with the supplied vapor species of the materials to

be synthesized when the substrate temperature is elevated beyond

their eutectic temperature. The alloy thus becomes a preferred

adsorption site for vapor since the liquid surface has a high

sticking coefficient. Nucleation occurs when the alloy is super-

saturated, and the axial growth proceeds by precipitation at the

liquid alloy/solid interface. Meanwhile, secondary nucleation

Nanoscale, 2012, 4, 1422–1435 | 1423


Fig. 1 A schematic of the VLS mechanism for the growth of ME (M ¼Zn, Cd; E ¼ O, S, Se, Te) nanowires. The growth consists of four stages:

(a) formation of metallic catalyst particles, either from intentional

introduction to the substrate or due to segregation of the catalyst film at

high temperature; (b) formation of alloy with the incorporation of the

supplied ME vapor into the catalyst; (c) supersaturation in the alloy,

resulting in the precipitation of a nucleation centre of ME crystal; (d)

axial elongation, culminating into ME nanowires. The schematic repre-

sents a randomly-oriented growth of nanowires, which does not neces-

sitate any chemical or epitaxial relation to the substrate.

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events are suppressed due to the higher energy involved than that

of crystal step growth, such that precipitation only occurs at the

single liquid/solid interface.17 1D growth continues until the

delivery of vapor reactant is discontinued or until the tempera-

ture is lowered below the eutectic temperature of the catalyst

alloy.

The presence of a metal-rich particle at the tip of the nanowire

is generally considered as an evidence of the VLS mechanism.16

However, the particle may also be detached due to interfacial

dewetting and thermal strain during the cooling process6 or

consumed during the growth.22 Noble metals (Ag, Pt, Ni, and

especially Au) are generally used as the catalyst particle, and an

appropriate choice of catalyst and growth temperature for the

synthesis of a certain II–VI compound can be determined with

the aid of pseudo-binary phase diagram.23 The catalyst particle

may also be used to control the diameter of the resulting nano-

wires, owing to the lateral growth confinement it provides.24

However, the metallic catalyst may diffuse from the tip22 and get

incorporated into the nanowire as a contaminant, which causes

concern, as such metallic contamination may affect the electronic

and optical properties of the nanowires.25

2.1.2. Self-catalytic growth. A special case of VLS growth

called self-catalytic growth arises when the constituent metallic

element of a compound serves as the catalyst, such as Zn for the

growth of zinc chalcogenides, thus eliminating the need for

a foreign catalyst which may evoke undesired contamination.26

To follow the mechanism, the catalyst can be generated by

pyrolysis of the source material at high vacuum and temperature

or through assisted decomposition by chemical reaction with the

carrier gas or additives. Similar to the conventional VLS mech-

anism, the catalyst particle may remain at the tip, as has been

demonstrated for ZnSe nanowire.27 However, the majority of

published works on self-catalytic growth reported the absence of

catalyst particle or any metal-rich site in the nanowires. Instead,

1424 | Nanoscale, 2012, 4, 1422–1435

tapered tip is often observed on the otherwise uniform diameter

nanowire.28–32 The lack of catalyst site can be ascribed to the

conversion of the catalyst into the material of the nanowire as

suggested from real-time TEM observation of the GaN nano-

wires growth.31 In absence of catalyst particle on the nanowire

product, the growth of nanowires without the use of foreign

metal catalyst can be attributed to self-catalytic growth by

consideration of the synthesis condition with regards to the

feasibility of catalyst generation and alloying with the vapor

reactant to follow VLS growth26–31 and the agreement of exper-

imental results with the kinetics and general characteristics of

a catalyzed growth.28 Improvements on the controllability of self-

catalytic mechanism, such as rational manipulation of the size

and positioning of the synthesized product, are still necessary to

fully exploit its demonstrated potential to produce high quality

and purity nanowires.

2.1.3. Vapor–solid (VS) growth. VS growth is a spontaneous

condensation process of vapor into solid material induced by the

decrease in Gibbs free energy from crystallization or the decrease

in supersaturation.12 VS growth is often invoked to explain the

growth of 1D nanostructures without any catalyst. Specifically,

the growth mechanism has been associated with the radial

thickening and tapering of nanowires, due to direct adsorption of

gas to the crystallized nanowires, and the growth of nanobelts.

Nanobelts33—also commonly called nanoribbons—are 1D

nanostructures that are subjected to a pronounced lateral growth

which accompanies the axial elongation, resulting in rectangular

cross sections. Nanobelts are generally conceived in the wurtzite

phase and numerous studies on nanobelts have been conve-

niently performed using II–VI materials, which can assume

a stable wurtzite phase under appropriate ambient conditions.34

Competition between surface energy minimization (which

favors the formation of nanobelts) and kinetics of crystal growth

(which controls the formation of nanowires) determines the

prevailing morphology during VS growth.35 When the vapor

precursor condensed at the substrate, the cation–anion adatoms

form a small nucleus while preserving the balance of local charge

and structural symmetry. Due to the difference in the kinetic

parameters for each crystal plane, anisotropic growth will result

and the nucleus will develop well-defined, low index crystallo-

graphic faces.9 However, when the temperature is sufficiently

high, accumulation of arriving atoms onto the smooth low-index

surface is prevented by the high diffusivity of the adatoms, thus

resulting in the expansion of surface area as more molecules stick

on the rougher growth front to form the nanobelts. Higher

supersaturation of the vapor phase has also been shown to

improve the 2D nucleation probability on the surface of

a nanowire to facilitate the growth of sheet-like structures,5,35

such that the dimension of a nanobelt can be controlled by

adjusting the growth temperature, the supersaturation ratio and

the growth time.

2.2. Experimental techniques

According to the setup and route to generate the vapor species,

the vapor based techniques commonly employed to synthesize

II–VI 1D nanostructures are thermal evaporation, laser-assisted



Fig. 2 (a) Typical experimental setup of the thermal evaporation

method to synthesize II–VI 1D nanostructures. Scanning electron

microscopy and high resolution transmission electron microscopy

(HRTEM) images of (b,c) ZnO nanowires, (d,e) CdSe nanowires, and

(f,g) ZnS nanobelts. The single crystallinity demonstrated by the uniform

fringes on the HRTEM images exemplified the high crystalline quality

that can be produced via the synthesis technique. (Fig. a,d,e are adapted

from ref. 28 with permission; Fig. b,c are adapted from ref. 36 with

permission; Fig. f,g are adapted from ref. 38 with permission.)

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growth, molecular beam epitaxy (MBE), and metal–organic

chemical vapor deposition (MOCVD).

The thermal evaporation technique is the most popular

approach by virtue of its simplicity and effectiveness in synthe-

sizing diverse nanostructure morphologies from various mate-

rials.29 The technique uses powder or condensed source materials

that are vaporized at an elevated temperature, where the result-

ing vapor phase is transported and deposited onto a substrate

within a tube furnace. The growth parameters are carefully

chosen to form the intended products: (a) the type of source,

which may determine the phase and stoichiometry of the

product; (b) the vaporization temperature of the source mate-

rials, by considering the volatility of the source; (c) the pressure

of the growth chamber, to control the evaporation rate and

vapor pressure of the source; (d) the substrate temperature,

which will strongly determine the type of product to be obtained;

(e) the carrier gas and its flow rate, to allow certain chemical

reaction to occur during vapor transportation and to tune the

growth rate; and (f) the evaporation time, which will influence

the amount and size of the products. Nanowires and nanobelts

from all of the zinc chalcogenides27,29,36–43 and cadmium chalco-

genides28,44–47 have been produced with thermal evaporation by

following the VLS and VS mechanism. A typical schematic of the

setup and selection of electron micrograph of 1D nanostructures

produced with the method are given in Fig. 2.

Laser ablation is a process whereby a laser, either pulsed or

continuous wave, ablates a solid target containing the desired

elements to construct 1D nanostructures.48 The clusters of

vaporized material are thus collected by a cold finger or a wafer

inside a growth chamber with a controllable environment, where

the nanostructures are formed. The synthesis with laser ablation

has since been accomplished with a broader range of material,

including ZnO,30 ZnS,49–51 ZnSe, CdS, CdSe23 and CdTe,52

demonstrating the general applicability of the technique.

MBE is an epitaxial growth technique of semiconductor

compounds involving the reaction of thermal molecular beams

on a substrate under ultra-high vacuum conditions.53 When the

level of a vacuum is sufficiently high (�10�8 Pa), the mean free

path of molecules may exceed the separation between the source

and the substrate. The condition, called the Knudsen regime,

allows the generation of the molecular beam to produce struc-

tures with great control of precision, stoichiometry and growth

rate. Several groups have demonstrated the applicability of the

approach for the growth of ZnO,54 ZnS,55 ZnSe56 and ZnTe57

nanowires. Nevertheless, the issue of low growth rate, impracti-

cality of ultra high vacuum, and the requirement of very high

purity sources have to be resolved before MBE can receive wider

usages in the industrial production of nanowires.

MOCVD is a deposition technique where at least one of the

precursor gases is a metallic atom (such as Zn from ZnO)

attached to an organic compound with a sufficiently high vapor

pressure. The precursor gas experiences pyrolytic reactions in

a furnace at elevated temperatures, where the metallic atom is

deposited on the substrate while the organic compound is

removed from the reaction chamber. MOCVD-grown nanowires

have also been demonstrated for II–VI nanowires.58–63 But as

opposed to MBE, MOCVD can accomplish high volume growth

with its high intake of precursor gas and possibility to use a large

diameter substrate to harvest more nanostructures. Yet,


difficulties exist for the MOCVD process, such as the use of

highly toxic gases and complicated chemistry, causing undue

increases of production cost.

2.3. Epitaxial growth of nanowire arrays

An unique scenario may occur when the II–VI nanowires are

grown epitaxially on the surface of the substrate, i.e., when the

crystal building block of the II–VI material is arranged in orderly

fashion by consistently following a specific orientational relation.

When the fast growth plane of the nanowires is the plane to be

heterointerfaced with the substrate, the nanowires will grow

perpendicularly from the substrate into a nanowire array. Unlike

non-heteroepitaxial nanowires which are randomly oriented

(Fig. 1 and Fig. 2a), epitaxial nanowire arrays possess strong

alignment in the vertical direction (such as Fig. 2b). Such align-

ment can be utilized to harvest certain orientation-dependent

Nanoscale, 2012, 4, 1422–1435 | 1425


Fig. 3 Recombination processes in a semiconductor. (a) Band-to-band

recombination. (b) Band-to-acceptor transition. (c) Donor-to-valence

transition. (d) Donor-to-acceptor-pair transition. (e) Recombination via

a deep center. (f) Non-radiative recombination via an intermediate state.

(g) Band-to-band Auger recombination.

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physical properties or phenomena exhibited by each nanowire,

which is essential in scalable applications.28 Additionally, in situ

growth in nanowire array structures also enables lattice relaxation

in lateral directions, allowing a larger critical height prior to defect

nucleations than that of similar lattice-mismatched epitaxial thin

films.64,65 Epitaxial nanowire arrays of II–VI materials were first

reported in ZnO,66 and have since been extensively studied in

a broad range of applications.

Conventional heteroepitaxy creates covalent bonds at the

heterointerface, thus requiring a matching of in-plane lattice

parameters between the nanowire and the substrate. Otherwise,

crystalline defects that are deleterious to the electronic and

photonic properties of the nanowires will occur. An epitaxial

strategy called van der Waals epitaxy has been proposed in thin

film growth67 to relieve the lattice matching requirements by

using substrates that promote van der Waals interaction—

instead of stronger chemical bonds—to connect the hetero-

junction. The van der Waals epitaxy has been attributed to the

growth of CdS, CdSe (Fig. 2a,b) and CdTe nanowire arrays with

excellent crystalline and luminescent qualities on a muscovite

mica substrate,28,68 thus also presenting a strong prospect for

application to other classes of compounds regardless of the

lattice mismatch.

3. Fundamentals of optical properties

3.1. Band structure

Optical transitions in semiconductors occur between the valence

and conduction bands. Therefore the studies of optical processes

and potential device applications require a complete knowledge

of the band structure. For II–VI semiconductors, the band

structure of wurtzite, zincblende, and rocksalt polytypes have

been theoretically investigated by varying the degree of

complexity. For wurtzite-type materials, the valence bands

originate from atomic p-like states whose symmetry is lower than

that of the diamond and zincblende cases. In comparison, the

minimum of the conduction band is s-like state with G7

symmetry.69 Spin–orbit coupling and crystal field split the

valence band into three sub-bands, resulting in a series of exciton

states which are denoted in the decreasing order of energy as A

(G9, also referred to as ‘‘heavy hole’’), B (G7, or ‘‘light hole’’), and

C (G7, or ‘‘split-off [SO] band’’).70,71 In order to observe different

excitonic transitions, one should consider the selection rules of

polarization.72 The transition energies of the intrinsic excitons

can be measured by employing low-temperature absorption,73

reflection,74 photoreflectance (PR)75 and photoluminescence (PL)

spectroscopy techniques.76,77 For example, the photon energy of

free excitons A and B (FX) in wurtzite ZnS nanowires was

determined by a low-temperature PL measurement to be 3.778

eV and 3.844 eV, respectively.78

3.2. Exciton and exciton complex

An exciton is a quasi-particle consisting of the bound state of an

electron and a hole interacting through the electrostatic

Coulombic force. Excitons behave similarly to a bosonic particle

freely moving in the crystal lattice of a semiconductor. In

comparison to free charge carriers, excitons demonstrate supe-

rior characteristics, such as large oscillator strength, narrow

1426 | Nanoscale, 2012, 4, 1422–1435

energy distribution of optical gain, and more varieties of possible

interactions with electrons or phonons. Moreover, excitons have

a dynamic electric dipole, which can couple with an electro-

magnetic wave, or light. In principle, this coupling between

a particle (exciton) and a field (electromagnetic cavity mode) can

be used to achieve zero threshold laser diodes.79 The exciton–

polariton coupling strength in the nanoscale is also significantly

enhanced from that in the bulk, as was recently reported with

CdS nanowire optical cavities.80 In 1D semiconductor materials,

we shall be concerned with Wannier excitons only.

There are two important optical processes which involve

emission of radiation from the semiconductors, namely, lumi-

nescence and inelastic scattering of light (also known as Raman

scattering). In a typical luminescent process, electrons in the

ground state are excited to a higher energy state, where after

some energy loss (i.e., relaxation) the excited electrons return to

the ground state by the emission of photons. There are several

categories of transitions that may happen in a PL process, as

shown in Fig. 3, depending on the sample and experimental

conditions such as temperature and excitation energy. For

instance, instead of a free carrier recombination processes, the

corresponding exciton recombination processes in the case of

a semiconductor with strong Coulombic interaction between the

electron and hole are FX, donor-bound excitons, and acceptor-

bound excitons.71 The donor or acceptor can be either neutral at

low temperature or ionized at high temperatures. The transition

energy difference between these excitonic emission and the free

carrier recombination processes is defined as the exciton binding

energy.

3.2.1. Exciton dynamics probed by steady state photo-lumi-

nescence. PL is used to study the optical properties of semi-

conductors. This is because PL provides very rich information on

both the intrinsic recombination processes and various radiative

recombination processes associated with imperfections of

a sample. With low temperature PL spectra, detailed information

about the semiconductor properties could be obtained, such as

exciton, phonon, and donor binding energy. Those properties are

very useful in giving a good picture of the optical quality of

semiconductor material.

Due to the difference in their densities of state (DOS), the

relative intensity of the defect and FX transition shows a very

different trend with the variation of excitation densities. As

shown in Fig. 4a, for ZnS nanowires at low temperature, the



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shallow level emission demonstrates a faster saturation

compared with FXs and could not be resolved at moderate

excitation density.78 It is necessary to note that the PL emission

for II–VI 1D nanostructures may show a gradual red shift with

increased excitation density, which can be ascribed to laser

heating effect commonly observed in other classes of

nanostructures.81

With the change of the temperature of the sample, a thermal

shift of the band gap of semiconductors can be observed. The

band gap shrinkage is related to the thermal dilation of the

crystal lattice as well as electron–phonon interactions. As shown

in Fig. 4b, the temperature dependence of the emission can be

described by both the Varshni formula82 and the Bose–Einstein

expression.83 Furthermore, the temperature dependence of full

width at half-maximum (FWHM) can be described by electron–

phonon coupling.84 Therefore, the determination of the exciton-

longitudinal optical (LO) phonon coupling strength is thus made

possible.

3.2.2. Optical absorption and ultrafast spectroscopy. Excitons

are generated from optical absorption, which are the case for II–

VI nanostructures with high exciton binding energies. The non-

equilibrium carriers will select channels (either radiative or non-

radiative) to return to equilibrium.85 Ultrafast spectroscopy—

such as time-resolved PL spectroscopy, optical pump–probe

absorption, transmission, reflection or second harmonic gener-

ation (SHG) spectroscopy—is a highly sensitive technique that

can be used to monitor the processes experienced by those

carriers, to measure the time required by the carriers to traverse

each channel (i.e., exciton lifetime), and to understand the

mechanism of the energy loss from the carriers (e.g. by exciton–

exciton, exciton–phonon, electron–exciton, or electron–phonon

interactions).

Exciton lifetime is very sensitive to crystal quality, where the

lifetime is longer for samples with improved crystal quality. We

have reported a bi-exponential decay for high-quality ZnS

nanowires grown by pulsed laser vaporization.78 Room-temper-

ature time-resolved photoluminescence (TRPL) measurements

produced decay times of 0.45 and 2.7 ns, most probably repre-

senting the non-radiative and the intrinsic radiative lifetime of

the free excitons, respectively.

Fig. 4 (a) Evolution of the PL spectra of ZnS nanowires measured at 10

K under different excitation densities. The dashed vertical lines are

a guide for the eyes. (b) Temperature-dependent PL mapping of ZnS

nanowires measured at an excitation density �9.4 mW cm�2. All data

were normalized by FXB intensity and the dashed lines follow the peak

positions. (Adapted from ref. 78 with permission.)


Interestingly, the ultrafast processes of carriers under one-

dimensional confinement change dramatically from their bulk

counterparts. For example, the lifetime of both electrons and

holes decrease with decreasing nanowire diameter, which are

attributed to the surface effects.86 Comparison between ZnO

nanowires and films also suggests that the decay lifetime of

electron in nanowires is half as long as that in films, which results

from the increased electron interaction with interfaces and grain

boundaries in nanowires.87

The following discussion focuses on the unique physical

properties of II–VI 1D nanostructures, especially those of

nanowires and nanobelts, for optical absorption and ultrafast

phenomena as compared to bulk materials and nanocrystals.

Firstly, nanowire absorption showed sizeable absorption cross

sections in visible region in the range of 10�12 cm2 mm�1 to 10�10

cm2 mm�1, which is roughly one order larger than those of

nanocrystals and is thus valuable for photodetector and photo-

voltaic applications.88 These 1D materials have intrinsically

anisotropic physical properties.89 As an example, polarization-

sensitive photodetectors can be fabricated from CdSe nano-

wires,90 and the origin of this polarization sensitivity stems from

the potential dielectric contrast effects as well as from the

confinement-induced optical selection rules. By PL spectroscopy

from single CdSe and CdSe/CdS core/shell nanowires, Giblin

et al. suggested that even though confinement may play a role in

determining the magnitude of observed anisotropy values,

dielectric contrast influences appear to dominate the overall

nanowire anisotropy response.91

Secondly, as compared to bulk materials, II–VI 1D nano-

structures have very rich luminescent properties induced by

intrinsic defects. For example, during vapor transport chemical

vapor deposition synthesis, stoichiometric intrinsic defects such

as such as vacancies (VS and VCd), interstitials (SI and SCd), and

anti-sites (SCd and CdS) for CdS for CdS nanobelts will be

unintentionally formed, as vapor pressures for II and VI group

elements are very different at certain temperatures. Fig. 5a shows

a PL spectrum measured at 10 K for CdS nanobelts and fitted

with multiple Gaussian functions to determine the peak positions

and width. The emission peak at 2.543 eV is attributed to neutral

donor bound excitons, historically labeled as I2. The peak near

2.531 eV is from I1, referring to excitons bound to a neutral

acceptor. A series of peaks between 2.416–2.266 eV demonstrate

characteristic multiple phonon replicas with an energy spacing of

approximately one LO phonon (i.e. 37 meV). Fig. 5b summarizes

the energy levels of all intrinsic defects of interest obtained by

first principle method. First principle calculations verify that the

shallow donors in CdS are contributed by sulfur vacancies while

the acceptors are contributed by cadmium vacancies with a spin

imbalance (Fig. 5c).92 Our results suggest that these two vacancy

states facilitate the formation of I2 with a fast decay dynamics on

the order of tens to hundreds of picoseconds, and a donor–

acceptor-pair (DAP) exciton complex with a much slower decay

process on the order of hundreds of nanoseconds due to the spin

polarization. Fig. 5d depicts the transient PL decay (normalized)

of I2 in CdS nanobelts at 10 K with pump power ranging from 1

mW to 20 mW.At low pump power, only one decay lifetime s1 dueto the single exciton decay was observed. As power increases

further, a faster decay lifetime s2 due to multi-exciton interac-

tions appears. Fig. 5e displays three emission decays centered at

Nanoscale, 2012, 4, 1422–1435 | 1427


Fig. 5 (a) PL fine structure of CdS nanobelts at 10 K with a 325 nm He-Cd laser excitation. The blue curves are Gaussian line-shape decompositions

with each peak clearly labeled. (b) Summary of the band diagram of CdS showing the intrinsic defect levels from the first-principle calculation (units in

eV). (c) Calculated electronic density of states (DOS) of CdS with cadmium vacancy (upper panel) and sulphur vacancy (lower panel). The arrows

indicate the defect levels, which also reveal a spin imbalance in the case of cadmium vacancy. (d) Transient PL decay (normalized) of I2 in CdS nanobelts

at 10 Kwith pump power ranging from 1 mW to 20 mW.At low pump power, only one decay lifetime s1 was observed. As power increases further, a faster

decay lifetime s2 due to multiexciton interactions was observed. (e) The normalized transient PL decay dynamics of the DAP emission at three different

wavelengths. (Adapted from ref. 92 with permission.)

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the three different wavelengths of the zero-phonon DAP PL

spectrum. Closer (more energetic) donor–acceptor pairs undergo

faster radiative recombination rates, which shows the charac-

teristic of DAP decay. Our results suggest the promise of engi-

neering the luminescent properties in terms of both energy and

lifetime of nanomaterials by controlling the species of defects.

Thirdly, the Auger process is an important non-radiative

channel, which happens frequently in nanostructures, especially

for the nanocrystals. The Auger recombination rates in 1D

systems are effectively suppressed, in contrast from those in

nanocrystals.93 The strongly reduced Auger decay rates lead to

increased optical gain-lifetime and hence efficient light amplifi-

cation, which will generate amplified spontaneous emission

(ASE) or lasing.94 This Auger process is population (pump

intensity) dependent in one-dimensional materials. Robel et al.

suggested that at high electron–hole pair densities, the dominant

process is bimolecular (exciton–exciton) Auger recombination,

while a three-carrier Auger relaxation mechanism which usually

happens in bulk materials and nanocrystals occurs under low

excitation intensity for small-diameter (between 6 and 8 nm)

CdSe nanowires.95

Fourthly, ASE or lasing in II–VI semiconductor nanowires or

nanobelts has been readily achieved with a threshold power

density ranging from several to tens of mJ cm�2 depending on the

quality and morphology of the sample.96 ASE takes place in the

absence of an optical cavity while lasing usually needs a cavity,

either formed as a natural cavity by the crystal facets or by

a mirror cavity, for further lasing from ASE process. The

ultrafast dynamics in the ASE (lasing) region are informative to

understand the fundamentals of nanophotonics. Time-resolved

1428 | Nanoscale, 2012, 4, 1422–1435

SHG and PL spectroscopy were first utilized to probe the

dynamics of single ZnO nanowire and nanobelt lasers.97 They

observed an 80 ps free exciton decay channel and a less than 10 ps

stimulated emission involving a quick channel. Recently, we also

observed the ASE process in CdS nanobelts, with stimulated

emission on the time scale of 20 ps.92 Dynamic competition

between I2 and DAP has been identified, which suggests that

compensation of acceptor levels in CdS nanobelts is required to

achieve ASE. This ASE process can also be influenced by the

band gap renormalization with wavelength-dependent lasing

time dynamics, showing different lifetime as the morphology of

the sample changes.96,98

3.3. Lasing

A variety of optically pumped stimulated emissions (SE) have

been observed by many researchers from II–VI semiconductors.

As discussed above, excitonic emission is preferable for

decreasing the threshold for lasing. In II–VI semiconductors, the

exciton binding energy of ZnO (�60 meV),76 ZnS (�40 meV),99

and CdS (�28 meV)100 is larger than the thermal energy at room

temperature, thus enabling the observation of excitonic emission

at room temperature. It is well known that the lasing action needs

a proper optical cavity to provide positive feedback. In the case

of one-dimensional nanowires, this is usually realized by either

pumping on a single wire (in Fabry–P�erot or whispering gallery

mode lasing) or observing random lasing in an ensemble of

densely distributed nanowires, both of which require not only

a proper pumping scheme but also a well controlled sample

manipulation or fabrication.



Fig. 7 The emission spectra fromCdSe NWs at room temperature under

different excitation densities. The inset shows the dependence of the PL

integrated intensity on the laser excitation density. (Adapted from ref. 68

with permission.)

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The near band emission of ZnO nanowires is located at �380

nm at room temperature. The first observation of lasing from

ZnO nanowire arrays was reported by Huang et al.66 The sample

was pumped byNd:YAG laser (266 nm, 3 ns pulse width) and the

emission was collected in the direction normal to the nanowires,

along the c-axis. As shown in Fig. 6a, when the excitation density

exceeds the threshold, sharp peaks with linewidths as small as 0.3

nm emerged in the emission spectra. The lasing cavity of these

nanowire arrays was ascribed to the natural cavity formed by the

two ends of the entire nanowires (Fig. 6c).

It is also possible that the laser cavity is formed by the entire

array, leading to lasing in arbitrary pathways with the nano-

wires as scattering centers. This phenomenon is known as

random lasing and was first observed in ZnO powder

samples.101 Random lasing in II–VI nanowires was previously

more commonly demonstrated in the UV regime with

ZnO.102,103 The random lasing in the near IR regime had not

been observed until very recently when high quality CdSe

nanowire arrays were achieved by the van der Waals epitaxy

mechanism. As shown in Fig. 7, near-infrared (NIR) random

lasing was observed at room temperature in CdSe nanowire

arrays grown on mica.68 Compared with conventional lasers,

random lasing from semiconductor NWs exhibits remarkable

differences. First of all, the random lasing can be observed

from all directions, and the spectrum recorded from a different

angle is different. Secondly, the lasing threshold strongly

depends on the excitation area. In the case of CdSe nanowires,

the pump area (Ath) and pump threshold (Pth) shows a A2/3th z

KP�1th relationship, where K is a constant.104,105 This indicates

that the vertically aligned CdSe nanowires should behave like

a quasi-3D random medium. Based on a similar concept,

random lasing in randomly-oriented nanobelts has also been

demonstrated in CdS.106

Fig. 6 (a) Emission spectra from nanowire arrays below (line a) and

above (line b and inset) the lasing threshold. The pump powers for these

spectra are 20, 100, and 150 kW cm�2, respectively. The spectra are offset

for easy comparison. (b) Integrated emission intensity from nanowires as

a function of optical pumping energy intensity. (c) Schematic illustration

of a nanowire as a resonance cavity with two naturally faceted hexagonal

end faces acting as reflecting mirrors. (Adapted from ref. 66 with

permission.)


3.4. Phonons

Phonons are quanta of lattice vibrations. They play a major role

in electrical, thermal, and optical properties of solid bulk mate-

rials and nanostructures107 as phonons carry energy and interact

with electrons via electron–phonon interaction and interact

among themselves via phonon–phonon interactions. The

importance of phonons are also demonstrated in various inter-

esting phenomena, such as phonon confinement in SNWs,108,109

phonon bottleneck effects in nanowires and quantum dot opto-

electronic devices,110 resonant Raman scattering,111 and deter-

mination of carriers mobility.107

Theoretically, the frequency and dispersion relation of

phonons can be obtained by solving the lattice dynamics equa-

tion. The vibrational modes of phonons can be either Raman-

active or IR-active. Hence, infrared absorption and Raman

spectroscopy are complementary in the experimental measure-

ment and assignment of vibrational modes.

Since the Raman effect was discovered more than 80 years

ago,112 Raman spectroscopy has become one of the most

powerful spectroscopic tools to study phonons and other element

excitations in condensed matter physics. Raman spectroscopy is

non-destructive and highly sensitive, where it can detect objects

on the order of a few nanometres, at a broad excitation energy

range (UV-NIR).

The first-order inelastic Raman scattering processes are two-

photon events that involve the simultaneous annihilation of an

incident photon and the creation of a scattered photon, mediated

by electron–photon and electron–phonon interactions.113 In this

process, the conservation of energy and momentum must be

satisfied simultaneously. In the visible and near infrared regime,

the momentum of the photon (wave vector) is less than 106 cm�1.

This wave vector is much smaller by about two orders of

magnitude than that of phonons, which corresponds to the

Brillouin zone boundary of typical crystals. Thus, the first-order

processes only access phonons and other elementary excitations

at or near the centre of the Brillouin zone, i.e., the well-known

q ¼ 0 Raman selection rule.

For one-dimensional nanowires, the shape of phonon disper-

sion relation can be obtained by folding the Brillouin zone of

bulk materials. However, some new modes (e.g. radial breathing

modes, RBMs) will emerge when the diameter of the wires is

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smaller than a certain critical size, apart from the notable

confined phonon states in the quantum size regime. Another

issue is the interface and surface vibrations, such as surface

phonons in nanostructures, which becomes more significant due

to the increase of the surface-to-volume ratio with the decrease of

size. The observation of surface optical phonons in nanowires

usually indicates some type of translational symmetry breaking,

since surface optical modes are not generally activated due to

momentum conservation.

3.4.1. Radial breathing modes (RBMs). The RBMs of

a nanowire are similar to the RBMs of carbon nanotubes

(CNTs),111 except that the vibration is considered for a solid

cylinder instead of a hollow tube. The mode corresponds to the

vibration of the atoms in the radial direction, as if the wire/tube

were breathing, where the frequency is strongly diameter

sensitive.

The frequency of RBMs based on elastic theory for a long

isotropic cylindrical nanowire, e.g., in silicon, is given by:114,115

u ¼ nl

R

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi8ð1� rÞð2� rÞ

s(1)

where n is the longitudinal sound velocity, R is the radius of the

nanowire, and r ¼ m/l + 2m, with (l,m) being the elastic constant

of materials. Eqn (1) indicates that the radial breathing mode

frequency follows a 1/R behavior, which is similar to the results

of RBMs in CNTs. Recent ab initio calculation shows the

frequency evolution of RBM can be fitted by the following

empirical power law for silicon nanowires:116,117

uRBMðdÞ ¼ 1

ARþ B(2)

whereA and B are the fitting parameters. Compared to elemental

semiconductor such as silicon, a more rapid progress in the

experimental study of RBMs in II–VI nanowires was witnessed

due to the wide availability of techniques to prepare nearly

monodisperse II–VI nanowires at a small diameter. For instance,

the presence of RBMs has been experimentally confirmed in

CdSe nanorods as shown in Fig. 8.118 The theoretical and

experimental results show that the RBM frequencies follow the

elastic theory analysis robustly, i.e. 1/R dependence, in spite of

the variation from cylindrical shape in the small diameter limit,

Fig. 8 Evolution of RBM frequency in CdSe nanorod and core–shell

structures as a function of the nanorod diameter. The theoretical calcu-

lation results and 1/R fitting are also given. (Fig. 8 is adapted from ref.

118 with permission.)

1430 | Nanoscale, 2012, 4, 1422–1435

and rather independently of the computational technique

adopted.

3.4.2. Surface optical modes. Surface phonon modes, which

are localized to the surfaces or propagating only along the

interfaces of a structure, are actually solutions of Maxwell’s

equations with appropriate boundary conditions.119 Surface

phonon modes are usually not Raman-active due to the

constraint of momentum conservation. They can be Raman-

active only if the translational symmetry of the surface is per-

turbed. Several ways have been tested to break such translational

symmetry, such as by surface roughness, by construction of

a grating along the surface, or by probing the evanescent wave

using a prism.120

Surface modes can now be optically probed in nanostructures

as the size modulation can break the translational symmetry

along a specific dimension, which provides the momentum in

order to activate surface modes in Raman scattering. Usually,

Raman scattering only probes optical modes as the momentum

of light is close to the wave vector of the optical mode in the

centre of the Brillouin zone, such that they are called surface

optical (SO) phonons.121,122 The SOmodes have been observed in

II–VI semiconductor nanostructures such as CdS,123 CdSe,124

and ZnS.125

For an infinite dielectric and semiconductor interface, the

surface mode frequency can be expressed by the following

equation:126

uSP ¼ uTO

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi30 þ 3m

3N þ 3m

r(3)

where 30 and 3N are the static and high frequency dielectric

constants, respectively, while uTO is the transverse optical (TO)

mode frequency at the zone center, and 3m is the dielectric

constant of the medium. The surface mode frequency is actually

an asymptotic limit of surface polaritons.126 Following eqn (3),

we can see that the SO frequency depends on the dielectric

constant of the surrounding medium. Therefore, SO phonons

can be confirmed by adjusting the 3m, or be used as a sensor to

specify the surrounding materials. Fig. 9a shows an experimental

evidence of SO mode in the Raman spectra from rectangular

cross-sectioned w-ZnS NWs carried out in three dielectric media.

In addition to a normal Raman band, an SO Raman band is also

activated and is seen to downshift as 3m is increased.125

In an infinitely long cylindrical nanowire and in the limit where

the phonon wave vector q[ u/c, the dispersion relation uSO for

a SO mode can be written as:121

u2SO ¼ u2

TO þ ~u2p

3N þ 3m f ðxÞ; x ¼ qr (4)

where, ~up is the screened ion plasma frequency, uLO is the LO

mode frequency at the zone center, and f(x) is the shape factor of

the nanowire. This model has been used to explain the experi-

mental observation of SO modes in Raman scattering in GaP

nanowires. For infinite diameter cylinders, i.e., r/N,f(x)/ 1,

then eqn (4) is equal to an infinite semiconductor flat surface

expressed by eqn (3).

For rectangular cross-sectioned nanowires, the anisotropic

boundary conditions result in two branches of SO phonon, i.e.,



Fig. 9 (a) The Raman spectra of w-ZnS nanowires collected in three

different dielectric media. The low frequency mode is the SO mode and

the high frequency one is the LO mode. (b) SO phonon dispersions for w-

ZnS nanowires with square cross-section in three different media.

(Adapted from ref. 125 with permission.)

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symmetric (S) and asymmetric (AS) SO phonons. Their disper-

sion relations are as follows:125

u2SOðqÞS ¼ u2

TO

30 tanhðqiLi=2Þ þ 3m

3N tanhðqiLi=2Þ þ 3m(5a)

u2SOðqÞAS ¼ u2

TO

30 cothðqiLi=2Þ þ 3m

3N cothðqiLi=2Þ þ 3m(5b)

For nanowires with an infinite edge length, i.e., L /

N,tanh(qL/2)/ 1,coth(qL/2)/ 1, both the symmetric (5a) and

asymmetric modes (5b) recover again to the surface phonons of

infinite flat semiconductor surface expressed by eqn(3).

Fig. 9b shows the SO phonon dispersion curves, for both

symmetric (S) and asymmetric (AS) modes, of a square cross-

sectioned w-ZnS nanowire in three different dielectric media 3m.

The S and ASmodes exhibit different dispersions at small qL, but

they converge to the same frequency as qL approach very large

values. Three horizontal dashed lines represent the experimental

SO phonon frequencies in the three dielectric media as shown in

Fig. 9b. The three experimental frequencies correspond to

approximately the same wave vector q. This is important self-

consistent evidence of the validity of dielectric continuum model

described in eqn (5a,b).

Fig. 10 Characterization of optical waveguiding in straight and bent

nanowires. (a) A scheme for scanning optical microscopy (SOM) illus-

4. Photonic and optoelectronic applications

II–VI 1D nanostructures offer a wealth of potential applications,

especially in photonics and optoelectronics due to their excellent

optical properties. This section highlights three of the interesting

instances of the applications: II–VI 1D nanostructures as wave-

guides, light emitting diodes, and photoconductors.
trating a focused laser spot scanned over the sample while monitoring
light emission from one end of the nanowire. The intensity at the end

indicated by the detector is plotted on a color scale as a function of laser

position to generate SOM images. (b) On the left, SOM image of a single

CdS nanowire; the reference end is indicated with a detector. Lighter

colors correspond to a greater end intensity. The scale bar is 10 mm. On

the right, dependence of the end intensity on the distance between the

laser spot and the end of the wire for a path that follows the wire.

(Adapted from ref. 129 with permission.)

4.1. Waveguides

In photonics, a waveguide is an object that is functionalized to

confine and channel electromagnetic waves into specific direc-

tions. Several inherent properties of II–VI 1D nanostructures are

immediately recognized to be suitable for waveguide applica-

tions. II–VI materials have wide band gaps and large dielectric


constants, yielding refractive index higher than their surround-

ings to allow transmission of light via total internal reflection

with a low loss. 1D nanostructures are also characterized by their

high aspect ratio, which effectively enables them to accomplish

their task as light interconnections.127

Various demonstrations of waveguiding using II–VI 1D

nanostructures have been accomplished since the surge of

interest in the nanowire field in the last decade. Johnson,

et al.128 reported the capability of ZnO nanowires in wave-

guiding, while simultaneously acting as optical cavities, anal-

ogous to a Fabry–Perot resonator, in the presence of flat and

parallel end faces on the nanowire. Barrelet, et al.129 showed

the possibility of CdS nanowires to act as active photonics

circuit elements for guiding light at sub-wavelength scales.

The nanowires are also observed to experience only moderate

loss of light propagation through sharp and acute bends, as

shown in Fig. 10. The finding allows further use of nanowire

waveguides in manipulating optical signals, as demonstrated

by Greytak, et al.130 for CdS nanowire electro-optic modula-

tors. Active waveguiding, when the nanowires are operated

near the band edge, was also demonstrated by the Zou group

for CdS nanoribbons131 and CdSe nanowires.132 Band edge

photoluminescence emission, excited within the body of the

nanowires, was transported along the wires to be emitted at

its end while no appreciable emission was observed in other

region.

The development in nanowire waveguides is anticipated to

pave the way for a highly integrated and compact photonics

circuit, which potentially mitigates the limitation of speed and

power dissipation in electronics to revolutionize communication

and computing systems. Thus, efficient injection, guiding, and

manipulation of light at sub-wavelength scale through 1D

nanostructures is evidently an essential research direction to

manifest the encouraging potential of II–VI 1D nanostructure

waveguides.

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4.2. Electroluminescence (EL) and light emitting diodes

(LEDs)

EL is a phenomenon in which electricity drives light emission,

including carrier injection and radiative recombination. EL has

attracted a great deal of attention due to the fact that it can be

easily integrated into electronic devices to provide wide appli-

cations including in lighting and displays. Studies are currently

being performed to resolve the main challenges in improving the

conversion efficiency and device stability.

II–VI semiconducting compound nanowires and nanobelts are

ideal candidates for EL because of their inherent properties such

as a direct band gap, large band gap coverage, high crystalline

quality, large surface area, and small size. A direct band gap is

the essential requirement for light emission while large band gap

coverage enables them to be utilized at various wavelengths from

IR to near UV. High crystalline quality can enhance the radiative

recombination rate while large surface area would provide the

flexibility to manipulate the surface. Besides, it is possible to

create a smaller beam spot due to the small size of the nanowires

and nanobelts.

Lieber and co-workers have realized electrically driven light

emission on n-CdS nanowire on a p-Si substrate.133 The Fabry–

Perot cavity modes formed inside nanowires lead to a series of

peaks at the emission spectra. Green light lasing can also be

obtained with a further increment of the bias, as shown in

Fig. 11. The heterostructure of n-CdS nanowire/p-Si nanowire

junctions also has been adopted to realize electrically driven light

emission.134 Additionally, multicolored light using the same p-

type nanowires can be achieved by changing the choice of n-type

nanowire. Hayden et al.135 utilized an n-CdS shell/p-Si core

structure to fabricate an LED. In addition, CdSe nanowires136

and nanobelts137 have also been made into devices to investigate

their respective EL properties. Besides, flexible single ZnO

nanowire light emitting diodes have also been demonstrated.138

4.3. Photoconductivity: photodetection and switching

When a beam of light shines on a semiconductor or an insulator,

the electrical conductance may either increase (positive photo-

conductivity) or decrease (negative photoconductivity). Great

efforts have been dedicated into this research field to find new

photoconducting materials with improved performance due to

their wide applications, such as in photodetectors, solar cells and

sensors.

Fig. 11 (a) Schematic of the n-CdS nanowire/p-Si device. (b) Electro-

luminescence spectra obtained from the device depicted in (a) with

injection currents of 120 mA (red) and 210 mA (green). (c) Electrolu-

minescence spectra from the device with injection currents of 200 mA

(red) and 280 mA (green) at 8 K. (Adapted from ref. 133 with

permission.)

1432 | Nanoscale, 2012, 4, 1422–1435

Utilizing a material in its nanoscale form has been shown to

improve greatly the photoconductivity in both carrier generation

and transport, which are the two processes involved in the study

of photoconductivity. The high crystalline quality of nano-

structures can reduce the scattering during the transport process

while the large surface-to-volume ratio can strengthen the

significance of surface states, which would enhance the recom-

bination rate and thus increase the response speed. In addition,

the large surface-to-volume ratio also improves the trapping of

charge carrier to enhance the sensitivity further, due to the more

pronounced depletion region near the surface.

Yang and co-workers initiated the research on photoconduc-

tivity of II–VI nanowires by studying ZnO,139 where the group

reported IV curves with and without illumination and the

switching characteristics as shown in Fig. 12. The large photo-

sensitivity and high on–off ratio demonstrates that nanowires are

good candidates for photoconductors. Since then, the photo-

conducting property of a wide variety of nanowires and nano-

belts—which is characterized by the sensitivity, spectral

response, response speed and switch characteristics—have been

studied in great detail, owing to the rapidly developing synthesis

techniques of 1D nanostructures. The photoconductivity of II–

IV 1D nanostructures such as that from ZnO,139,140 CdS,141–144

CdSe,90 CdTe, ZnSe,145 and ZnS146,147 have been intensively

investigated. Various fascinating phenomena on the photo-

conducting properties of II–VI 1D nanostructures have also been

studied, including the improvement of photosensitivity by

exploiting a Fabry–Perot cavity between a nanowire or nanobelt

photodetector with the substrate90 and a polarization-dependent

sensitivity, as was observed in CdSe90 and CdTe nanowires148 due

to the high aspect ratio of nanowires.

5. Conclusions

We have provided an overview of the material preparation

techniques, the fundamentals of exceptional optical properties

and the promising optoelectronic applications of II–VI 1D

nanostructures, along with brief highlights of prominent past

research activities in the field. With the rapid progress of the

research in the field within only the last decade, it is thus

Fig. 12 (a) I–V curves of a single ZnO nanowire under 365 nm, 0.3 mW

cm�2 UV light illumination, exhibiting the dark and photocurrent

property of the structure. The inset shows an FE-SEM image of a 60 nm

ZnO nanowire bridging four Au electrodes. (b) Reversible switching of

a ZnO nanowire between low and high conductivity states when the

handheld UV-lamp was turned on and off. The bias on the nanowire is

1 V. (Adapted from ref. 139 with permission.)



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increasingly important to consider the marketability of the

promising technological products that II–VI 1D nanostructures

can offer. As an important bottom-line, novel technology must

be able to achieve lower cost-per-performance than the current

equivalent established technology before it can be fully utilized.

Detailed comprehension of material properties and device

behavior of II–VI 1D nanostructures are needed, due to their

relevance on both to the performance reliability of the produced

technology and to the emergence of revolutionary applications

from the novel fundamental properties manifested at the nano-

scale. In particular, we believe that more intensive research

studies should be devoted to gaining deeper understanding of the

properties of excitons and phonons within the II–VI 1D nano-

structures and their coupling with other quanta such as photons

and plasmons, as has been demonstrated to strongly affect the

overall optical and emissive properties of the materials.80,149

Several issues in the field of II–VI 1D nanostructures also need

to be resolved to actualize a real world impact. Improvements in

the preparation techniques are still necessary to achieve more

control with simultaneous precision on the composition, physical

properties, dimension, location, alignment, and orientation of

the nanostructures for massive-scale device integration. Addi-

tionally, it is also important to perform studies investigating

‘‘better or worse’’ questions and to address important side issues,

such as device endurance and toxicology,150 to support the

abundant proof-of-concept oriented studies that are still being

the paradigm of the current state of research on II–VI 1D

nanostructures.

Acknowledgements

Q. X. acknowledges the strong support from the Singapore

National Research Foundation (NRF-RF2009-06), generous

start-up grant support (M58110061) and the New Initiative Fund

(M58110100) from Nanyang Technological University.

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