Nanoparticles and nanowires: synchrotron spectroscopy...

1194 Int. J. Nanotechnol., Vol. 5, Nos. 9/10/11/12, 2008

Copyright © 2008 Inderscience Enterprises Ltd.

Nanoparticles and nanowires: synchrotron spectroscopy studies

T.K. Sham Department of Chemistry, University of Western Ontario, 1151 Richmond Street, London Ontario, N6A 5B7 Canada Fax: 1-519-661-3022 E-mail: [email protected]

Abstract: This paper reviews the research in nanomaterials conducted in our laboratory in the last decade using conventional and synchrotron radiation techniques. While preparative and conventional characterisation techniques are described, emphasis is placed on the analysis of nanomaterials using synchrotron radiation. Materials of primary interests are metal nanoparticles and semiconductor nanowires and naonribbons. Synchrotron techniques based on absorption spectroscopy such as X-ray absorption fine structures (XAFS), which includes X-ray absorption near edge structures (XANES) and extended X-ray absorption fine structures (EXFAS), and de-excitation spectroscopy, including X-ray excited optical luminescence (XEOL), time-resolved X-ray excited optical luminescence (TRXEOL) and X-ray emission spectroscopy (XES) are described. We show that the tunability, brightness, polarisation and time structure of synchrotron radiation are providing unprecedented capabilities for nanomaterials analysis. Synchrotron studies of prototype systems such as gold nanoparticles, 1-D nanowires of group IV materials, C, Si and Ge as well as nanodiamond, and compound semiconductors, ZnS, CdS, ZnO and related materials are used to illustrate the power and unique capabilities of synchrotron spectroscopy in the characterisation of local structure, electronic structure and optical properties of nanomaterials.

Keywords: gold nanoparticles; semiconductor 1-D nanowires; synchrotron radiation; XAFS; XANES; XEOL; TRXEOL; electronic structure and optical properties.

Reference to this paper should be made as follows: Sham, T.K. (2008) ‘Nanoparticles and nanowires: synchrotron spectroscopy studies’, Int. J. Nanotechnol., Vol. 5, Nos. 9/10/11/12, pp.1194–1246.

Biographical notes: T.K. Sham received his PhD in Chemistry from the University of Western Ontario in 1975. Earlier, he received his BSc (Hon. 1st class) from the Chinese University of Hong Kong. After a Postdoctoral Fellowship at UWO, he was on the staff of the Chemistry Department at Brookhaven National Laboratory before returning to UWO in 1988. He is presently Professor of Chemistry and Canada Research Chair in Materials and Synchrotron Radiation. Since 1999, he has been the Scientific Director of the Canadian Synchrotron Radiation Facility (CSRF) at the Synchrotron Radiation Center, University of Wisconsin-Madison. He is also a beam team leader at the Canadian Light Source. His research interests are studies of the interplay of electronic structure, materials properties and spectroscopy using a variety of techniques, especially synchrotron radiation. Emphases are placed on materials in unusual dimensions, such as surface, interface and nanostructure, and the development and application of synchrotron radiation.

Nanoparticles and nanowires: synchrotron spectroscopy studies 1195

1 Introduction

The scope of this paper is to review several areas of nanostructure research that have been conducted in our laboratory over the last decade. This choice to focus on our own work is in accord with the theme of this volume. The field of nano research has seen enormous growth in the last decade and its breath and anticipated impact are arguably unprecedented in the history of science and technology. Our research concentrates on the interplay of nanomaterial properties, electronic structure and characterisation techniques, especially the use of synchrotron radiation spectroscopy and its implications. The development of synchrotron radiation technology has been an important element in our research. It happens that synchrotron radiation technology has made great strides temporally parallel to the growth of nanoscience and nanotechnology, providing a variety of powerful characterisation tools not available in conventional laboratories for research in nanoscience and nanotechnology. Although this review is based primarily on our work, which has a strong synchrotron flavour and deals with the synergy between nanomaterial research and synchrotron techniques, the results will be presented in the context of contemporary research.

Nano research means this to one but that to the other. However, what is of common interest is the effect of the size of materials, of the order of nanometres (nm) to tens of nanometres, on their properties. In the context of our discussion, we will limit ourselves to the study of the structure, morphology and electronic behaviour of nanosystems of metals and semiconductors that are physically confined in one, two or three dimensions to the order of nanometers or tens of nanometres, especially nanoparticles and nanowires. Under these confined conditions, materials can behave very differently in their electronic, magnetic, mechanical, chemical and optical properties from that of their bulk counterparts. Depending on the dimensionality of the confinement, these materials are often called quantum wells (1–D confinement), quantum/nano wires (2-D confinement) and quantum/nano dots (3-D confinement). There is also a combination of the above, such as quantum sponge (e.g., porous silicon) and many variations such as rods, onions, tubes, needles, ribbons, rings, zigzag wires, etc.

Dramatic changes of properties of nanomaterials from that of the bulk when the size of a material is in the dimension of nanometers owes to the wave behaviour of the electron and its de Broglie wavelength. Electrons in matter will behave differently from their bulk when they are confined in the dimension of the order of nanometers where the size of confinement is comparable to the wavelength of the electron. The simplest illustration of the effect of this change of boundary conditions is the widening of the separation of energy levels when we reduce the length of the potential well in a one-dimensional particle-in-a-box problem. This illustration has been widely and successfully used, albeit qualitatively, to explain the band gap opening of semiconductor nanostructures and its size dependent optical properties (e.g., blue shift of the near band-gap emission in semiconductor quantum dots). One of the earliest examples of a nanostructure exhibiting dramatic property change in last 15 years or so is porous silicon, a network of pillars and nodules of crystalline silicon of the size of a few nanometers left behind after electrochemical etching of a silicon wafer in HF solution [1]. This approach, like lithography, is now commonly referred to as the top-down method. Remarkably, porous silicon emits visible light at room temperature upon photo or electrical excitation [2]. Most interestingly, the wavelength of the luminescence can be tuned by varying the size of the nodules and pillars within the nm regime – the smaller

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the size, the shorter the wavelength of emission, and chemical modification of the surface [1]. Bulk silicon, an indirect band gap material, does not emit light in the visible.

There are other implications when a material is in the nano dimension, most importantly is the increase in surface to bulk atom ratio, which modifies the surface chemistry (relevant to catalysis) and the transport properties of charge carriers. The latter involves the scattering of the electrons off the inner surface (wall) of the nanowires; thus gas adsorbed on the surface can modify the conductivity markedly (sensors applications). The large surface to bulk atom ratio can alter the melting point since surface atoms have lower cohesion energy, etc.

We will describe here relevant synchrotron spectroscopies and the information they can provide in case studies of several prototype nanosystems using conventional and synchrotron techniques. This paper is organised as follows: We first describe briefly in Section 2 the nano systems of interest together with preparation strategies. This is followed in Section 3 by the description of characterisation using conventional and synchrotron methodologies relevant to the analysis of nanomaterials. After that, representative cases will be discussed in Section 4 with the emphasis on the role of synchrotron radiation techniques in characterisation. The prospects and the future role of synchrotron radiation in nano research are discussed in Section 5.

2 Nano systems and relevant issues

We shall describe below representative nano systems prepared in our laboratory although variations of these systems have been prepared elsewhere with similar or different strategies. Some general preparation techniques are noted. The systems that we will describe here are

• Si and Ge nanostructures

• Au nano particles and related systems

• compound semiconductor nanostructures (CdS, ZnS, ZnO) and related systems

• carbon nanostructures (nanodiamond and nanotube) and related systems.

The challenges of nano research are the control of shape and size in fabrication, the control of proximity effects (interaction of nanostructures with the surrounding and among themselves) and the control of assembly of nanostructures, the ultimate requirement for nanotechnology. The shape and size control is the starting point if one is to understand size-dependent properties. Most of our studies described below are based on systems with a well-defined range of sizes. It is useful to describe briefly our synthesis strategies for the fabrication of these systems, which contain general and some unique features. Except for the preparation of porous silicon, which was based on a top-down method, all the strategies we have considered are bottom up methods (from atoms). Let us first consider the preparation of semiconductor nanowires [3–6]. We utilise a thermal evaporation method with a tube furnace, in which the starting materials, often metal oxide powder, SiO or GeO, for example, are heated in one side of the tube and transported down stream by a carrier gas. Nanostructures are deposited onto a pre-selected substrate, which is sometimes coated with a catalyst (e.g., Au nanoparticles). This method is based on the VLS (v-liquid-solid) mechanism [7] with some variations


(see below). This technique takes advantage of the disproportionation of the oxide and the formation of eutectic of the catalyst with the material (e.g., Ge) of interest in liquid. It should be noted that except for Si nanowire (SiNW), the preparation of all the nanowires discussed here required the use of the catalyst. Figure 1 shows the phase diagram of Au-Ge [8] and the resulting Ge nanowire (GeNW) [9–12]. The analysis of the results will be discussed later.

Figure 1 (a) Phase diagram for the Au-Ge system showing phase separation as the system cools from 800°C and (b) TEM of Ge nanowire showing the Au particle of 40–50 nmat the tip

(a) (b)

The nanoparticle (NP) systems to be discussed here are Au and Rh NP. Au NPs were prepared with both hard-template and soft template techniques [13–18]. The soft template technique is a modification of conventional colloidal chemistry in which metal ions are reduced in solution in the presence of a capping molecule, which controls the rate of aggregation of the metal atoms (reduced from ion), hence the size of the NP. The chemistry of the capping molecules can also be tuned from weakly interacting species such as the -OH and –NH2 groups in a dendrimer, to strong interacting species such as thiol, R-SH, yielding nanoparticles with different properties albeit with the same size.

Figure 2 shows the schematic for capped NPs and representative weak and strong capping molecules. The hard template technique takes advantage of the reducing properties of the functional group on the surface of a substrate, which itself is a nanostructure. For example, HF treated porous silicon and silicon nanowire [1,13,18] both have a hydrogenated surface, which serves as a reducing agent for the reduction of metal ions from solution by an immersion method. This reductive deposition technique can be further modified using a combination of soft template, hard template, and electrochemistry. For example, Au NPs have been deposited electrochemically onto a silicon cathode in the presence of capping molecules. The difference between deposition with and without capping ligand is dramatic as displayed in Figure 3.

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Figure 2 Schematic (from left to right) for dendrimer (top left) and thiol (bottom left) capped nanoparticle and dendrimer with amino group and thiol ligands

Figure 3 Morphology of electrochemically reduced Au nanoparticles on a Si substratein the absence (a) and presence (b) of a capping molecule, thiol

3 Conventional and synchrotron methodology for nanosystem analysis

Important issues such as morphology, crystal structure, electronic structure, optical properties, chemical properties, etc., are to be addressed in the analysis of nanosystems. Many conventional and synchrotron techniques have been used in the studies of the systems described below. This section describes these techniques and their applicability in nanoanalysis. Emphasis is placed on synchrotron radiation.

3.1 Conventional techniques

The conventional imaging techniques used in the analysis include Scanning Electron Microscopy (SEM), Transmission Electron Microscopy (TEM) and High Resolution Transmission Electron Microscopy (HRTEM). These tools not only can reveal the morphology of the nanostructure, sometimes with atomic resolution, but also are equipped with energy dispersive X-rays (EDX) analysis, selected area electron diffraction (SAED) and electron energy loss spectroscopy (EELS) capabilities. Another imaging technique, the scanning probe microscopy (SPM) based on Atomic Force Microscopy (AFM) and Scanning Electron Microscopy (STM), has also been used. The structure of


nanomaterials is often characterised with X-ray powder diffraction (XRD), which can reveal the crystal structure and the size of the crystallite via the analysis of the width of the diffraction peaks. UV-visible spectroscopy has been used to detect plasmon resonance of metallic nanosystems and to perform luminescence studies. X-ray photoelectron spectroscopy (XPS) has been used to study nanomaterials through the analysis of the core level binding energy shift and the valence band.

3.2 Synchrotron radiation spectroscopy

Synchrotron radiation (SR) is highly collimated electromagnetic radiation emitted tangent to the orbit (storage ring) where electrons are circulating at a speed close to the speed of light [19,20]. SR is emitted when the electron is bent (centrifugal acceleration) by a magnetic field. SR is brighter then the sun and is tunable from IR to hard X-rays. The useful range of SR from a storage ring depends on the energy of the electrons circulating in the orbit and the magnetic devices used to bend them. In general, the larger the electron energy is, the harder (more energetic) the X-rays. Modern multi-purpose synchrotrons are known as third generation light sources. Such synchrotrons use insertion devices called undulators and wigglers to perform many mini bends on the electrons with weak and strong magnetic fields, respectively, when electrons are coming through the insertion devices installed in the straight section. These devices increase the brightness of SR and select a desired wavelength region for specific experiments. A medium energy third generation storage ring like the Canadian Light Source (CLS) operates at 2.9 GeV with a current of 500 mA.

The CLS accommodates experimental stations that can be used to conduct a variety of experiments taking advantage of the tunable bright light. It is the brightness and the tunability of the photons with a continuous energy range from infrared to UV, to VUV (tens of eV) to soft X-ray (hundred to a few thousands of eV) and to hard X-ray (several thousands to tens of thousands of eV) that provide unprecedented capabilities for probing the structure and electronic properties of matter using absorption spectroscopy and related techniques. This wide range of photon energy can access practically all element-specific core levels. In addition, synchrotron light is also highly polarised (linearly in the plane and can be made elliptically polarised with a special insertion device) and has a desirable time structure for texture and timing experiments, respectively. The electrons stored in the ring circulate in short bunches, which are often equally spaced. Thus every time the electron bunch comes by the experimental port, it emits a light pulse, typically of the order of 100 ps (pulse width) with a repetition rate (separation between pulses) of tens to hundreds of ns depending on the size of the ring and the number of bunches that are filled.

The SR techniques for nanostructure analysis described here are:

• XAFS (X-ray Absorption Fine Structures) which is often divided into the study of the near edge region (up to ~50 eV above the threshold), often referred to as XANES (X-ray Absorption Near Edge Structure) or NEXAFS (Near Edges X-ray Absorption Fine Structures), and the extended energy region, EXAFS (Extended X-ray Absorption Fine Structures), the study of the oscillatory modulation of the absorption coefficient beyond the XANES/NEXAFS region to as much as 1000 eV above the threshold. Although this division of the two regions is somewhat arbitrary, it does reflect the dominant scattering behaviour of low kinetic energy photoelectron

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(multiple scattering) and high kinetic energy photoelectron (single scattering) by the atomic environment in the near edge and extended energy region above the threshold, respectively.

• XEOL (X-ray Excited Optical Luminescence) in both energy (wavelength scan) and time domain (using the synchrotron pulse as a reference), as well as time-gated spectroscopy by selecting a time window, and to a lesser extent.

• XES (X-ray Emission Spectroscopy) in which, the fluorescence X-rays were energy analysed. Selected synchrotron spectroscopy techniques discussed in this review are summarised in Table 1.

Table 1 Selected synchrotron spectroscopy techniques in nanomaterials analysis

Technique Description Information XAFS(XANES/NEXAFS)

Measurement of the X-ray absorption coefficient as a function of photon energy from below to ~50 eV above an absorption edge of a core level of an element in a chemical environment

Local symmetry and bonding – unoccupied densities of states

XAFS (EXAFS) Measurement of the modulation of the absorption coefficient as a function of photon energy beyond the XANES region to as much as 1000 eV abovethe edge

Short range order around the absorbing atom – coordination number, bond length andlocal disorder (Debye – Waller factor)

XPS X-ray (tunable) in, photoelectrons -out, the kinetic energy of the photoelectrons is analysed

Core level binding energy and valence band – chemical information

XEOL/TRXEOL X-rays in, optical photons – out (Near IR, visible and UV), can be measured in energy (wavelength scan) and time domain (decay lifetime using the synchrotron pulse as a reference)

Origin of luminescence from light emitting materials (element, site, excitation channel and time specific)

XES X-ray photons – in, X-ray photons (fluorescence) out, the fluorescence X-ray is analysed (monochromator)

For shallow cores below the valence band, XES probes the densities of states below the Fermi level (valence band)

3.2.1 X-ray Absorption Fine Structures (XAFS) spectroscopy

XAFS refers to the modulation of the absorption coefficient above an absorption edge of an element in a chemical environment compared to that of a free atom [21–23]. It refers to the details of the X-ray absorption (XAS) spectrum above an absorption edge (threshold for the excitation of a core electron) of an element. XAFS arises from dipole transitions from core to unoccupied states governed by the molecular potential of the environment of the absorbing atom. XAFS is conveniently divided into the XANES region (up to ~50 eV above the threshold), sometimes known as NEXAFS [24] and the EXAFS [21–23], which extends to as much as 1000 eV above the edge. This division at ~50 eV is somewhat arbitrary as noted above, but it does reflect the dominating physical processes involving the photo excited electron (from multiple scattering in XANES to single scattering in EXAFS) and the corresponding information it can provide.


In the XANES region [24,25], dipole allowed excitation (e.g., 1s to 2p, at the K-edge) turns on when the excitation energy approaches the threshold energy required to create a core hole (1s hole for carbon for example). As the excitation energy increases across the threshold, the core electron is excited to previously unoccupied electronic states that are bound (e.g., LUMO in molecules, bottom of the conduction band in semiconductors and insulators, and the unoccupied states above the Fermi level in metals), quasi-bound (shape resonance, or virtual orbitals, multiple-scattering states or higher energy bands in solids) and then in the continuum (with kinetic energy >0 eV). The photoelectron produced in the XANES region has low kinetic energy and is subjected to multiple scattering by the surrounding atoms. Thus resonance in the XANES can be treated as bound-to-bound and bound to quasi-bound transitions in molecule or transitions to the unoccupied states above the Fermi level in solids. Therefore, XANES are sensitive to the unoccupied densities of states (e.g., population of the LUMO etc.) and local symmetry (selection rules).

In the EXAFS region, the kinetic energy of the photoelectron increases with increasing excitation energy such that single scattering becomes dominant and the modulation of the absorption coefficient (µ), the EXAFS, χ(k) = [µ(sample) – µ(free atom)]/µ(free atom) can be treated with standard EXAFS formalism, χ(k) = A(k)Φ(k), containing an amplitude A(k) and a phase term Φ(k), where k in Å–1 can be expressed as

12

(eV) (eV) (eV) (eV)(Å )

/ 2 3.81(eVÅ )o oE E E E

km

− − −= = (1)

22( ) ( ) ( , 2 ) exp[ 2 ]exp[ 2 / ]ii i i

i

NA k S k f k k RkR

π σ λ= − − (2a)

( ) sin[2 ( )]i ik kR kφΦ = + (2b)

where E is the photon energy and Eo is the threshold energy. A(k) contains all the information about the magnitude of the scattering (size of the modulation). It has three important terms that are chemically sensitive. These are the backscattering amplitude of the neighbouring atom i, fi(k, 2π), the coordination number of neighbouring atoms of the same kind, Ni, and a damping term, the Debye-Waller factor, exp[–2k2σi

2], which describes the effect of the root mean square displacement, σi, of the interatomic distance Ri between the absorber and the scatterer (from thermal effect and static disorder) on the amplitude. There are also the slow varying many body term, S(k), and the kinetic energy dependent energy loss terms associated with the mean free path of the photoelectron λ.For nano materials, the coordination number and the Debye-Waller are very sensitive terms because, f(k, 2π) remains the same for bulk and nano material but the effective Ni becomes smaller as the size of the nanostructure becomes smaller. This is often accompanied by static disorder (slightly different bond lengths due to surface relaxation and reconstruction). The theory is now reasonably developed for evaluating the phase and amplitude functions [26]. To illustrate these effects qualitatively, we compare below the Si K-edge XAFS of the bulk, nanowire and porous silicon [27].

The phase term Φ(k) contains the information about the interatomic distance as well as the phase shift of the absorbing atom and the backscattering atom, φ(k). It has a sinusoidal dependence in k space. This accounts for the oscillation of the absorption

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coefficient in k space. The amplitude and the phase term can be separated by Fourier Transform or by curve fitting yielding information about the inter-atomic distance, coordination number and local disorder. The effect of the phase can also be seen in Figure 4. In general, for chemically similar systems, the shorter the bond is, the farther apart the separation of the oscillation maximum is. This can be illustrated with a linear approximation of φ(k) in k space, or φ(k) = α + βk where α and β are constants [28]. Thus between the adjacent EXAFS maximum, ∂Φ(k)/∂k = 0, or 2kRi + α + βk = nπ/2, n = 1, 3, 5 etc., we have

∆kmax = π/(2Ri + β) or Ri = [π/∆kmax – β]/2. (3)

We can see from the above equation that the separation of the adjacent maxima of the EXAFS in k space, ∆kmax, is inversely proportional to the bond length. This means that the larger the progressive separation, the shorter the bond. We will return to this below. The theory of EXAFS has been well established and will not be described here. In the analysis of Au nanoparticle EXAFS discussed below, we use the Fourier transform (FT), which separates the amplitude and the phase function, as well as fitting to compare the inter-atomic distance in these nanostructures as a function of particle size and surface chemistry. For a qualitative illustration, we show in Figure 4 the Si K-edge XAFS of Si(100) (most order), Si nanowire (less order) and porous silicon (PS) (least order).

Figure 4 Si K-edge of Si(100), SiNW and PS showing the XANES and the EXAFS. The vertical lines align the maximum of the oscillations

It should be noted that the y-axis of Figure 4 is in TEY (total electron yield). In XANES/NEXAFS measurements of low Z (atomic number) elements, direct transition measurement is not always practical since the absorption cross-sections is so large that a very thin uniform specimen is required. For example, the 1/e attenuation length of X-rays at energy above and below the carbon K-edge (~290 eV) is 2.4 µm and 0.089 µm, respectively. Thus, the measurement of the absorption coefficient is often carried out with total electron yield (TEY), X-ray fluorescence yield (FLY) and for light emitting materials, photoluminescence yield (PLY) in lieu of transmission. It is assumed that the yield is proportional to the absorption coefficient. This assumption is valid for TEY and may suffer from self-absorption (broadening of the peaks) in FLY and PLY measurements. Some of these issues will be addressed later.

A couple of interesting features from Figure 4 can be used to illustrate what one can observe in XAFS as the crystal size becomes small, e.g., from bulk Si in Si(100) to nanostructure in SiNW and PS. First, all spectral features are very similar with the XAFS oscillations appearing in the same energy for the three samples, indicating that they are


all elemental Si with identical diamond-like crystal structures. Second, the spectral features, which are sharp, in Si(100) become blurry in SiNW and blurrier still in PS; this is accompanied by a reduction in amplitude. This observation indicates that there is a reduction in the amplitude function A(k) due to degradation in long range order since the crystallites are much smaller (order of nm to 10 nm) in SiNW and smaller still in PS (several nm). Both disorder and reduction in coordination number contribute to the reduction in the overall amplitude of the EXAFS, which are, in most cases, apparent by inspection of the EXAFS region (modulation of the absorption coefficient). The EXAFS for Si(100) and SiNW in k space and corresponding FT are shown in Figure 5, confirming amplitude reduction observed in the absorption spectra [27].

Figure 5 Amplitude function (left) and its Fourier Transform (right) showing the phase uncorrected distribution of the first and higher shell Si-Si inter-atomic distance in Si(100) and SiNW

3.2.2 XEOL in the energy and time domain

XEOL is a photon-in photon-out technique in which the system is excited with X-rays from a synchrotron light source via the excitation of primarily core electrons [29]. The decay of the core hole (via Auger, dominant for elements of low atomic number, and X-ray fluorescence) of an atom in a solid, e.g., a semiconductor, ultimately leaves the excited system with holes in the valence band (HOMO and HOMO –1 etc., in molecules) and electrons in the conduction band (LUMO, LUMO + etc.). The radiative recombination of these electrons and holes produces luminescence. The presence of surface and bulk defects (energy levels within the band gap) of semiconductor or inorganic solids (e.g., metal oxide) is also responsible for the luminescence. The advantage of XEOL with synchrotron radiation over conventional excitation sources (lamp and laser in the visible and UV region) is that tunable X-ray photons allow for the preferential excitation of a particular element (e.g., C in OLED materials) and a particular excitation channel at an absorption edge (1s – π* resonance associated with the aromatic ring, σ* associated with C-H bond, etc., at the C K-edge) since core excitation thresholds are element specific. Thus, with high photon energy resolution that is readily obtainable with modern synchrotron light sources and monochromators, XEOL can be site (element), chemical (oxidation state), and excitation channel specific [30,31]. A schematic for the XEOL and XES technique is given in Figure 6, which illustrates the origin of light emission from light emitting semiconductor or insulator as well as the valence densities of states of the valence band from XES.

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Figure 6 Schematic for XEOL and XES. The left panel shows the creation of the core-holeby photoabsorption. The right panel (top) shows the electrons and holes in conduction and valence band, respectively resulting from the relaxation of the excited state as well as defect levels in the band gap. The radiative recombination yields hvop. The right panel (bottom) shows the X-ray fluorescence photon hvf (for colours see online version)

The luminescence intensity at a given excitation photon energy, PLY, either in zero order (total intensity) or at a selected wavelength, has been used to monitor the absorption across an edge. This measurement is called optical XAFS, and can be used to reveal the origin of the luminescence. This is especially true with soft X-ray excitation. The details of de-excitation in the optical region are complex since the system is in a super-excited state after X-ray excitation. In general, excited states with a core-hole decay via X-ray fluorescence and Auger, followed rapidly by a phonon/internal conversion route or intersystem crossing to the luminescence chromophore (triplet in organometallic phosphors [32], for example). This is why XEOL normally exhibits similar spectral features independent of excitation energy. The PLY however is very sensitive to the effective coupling of the excited state to the luminescence chromophore (efficiency in producing electron-hole pairs that recombine radiatively) and the history of the secondary processes (sample thickness). In the case of composites or heterostructures where more than one site is luminescent and all sites are competing for the excitation photons across an absorption edge, discernible excitation energy-dependent XEOL and optical yield dependent XANES can be observed readily [30]. Figure 7 shows the XEOL from porous silicon excited at the Si and Si oxide resonance at the Si K-edge.

XEOL in the time domain or TRXEOL (Time-Resolved X-ray Excited Optical Luminescence) takes advantage of the short synchrotron pulse and the repetition rate using multi-bunch and single bunch mode [32,33]. At the Advanced Photon Source (APS) at Argonne National Laboratory, the normal top-up mode provides a synchrotron pulse of ~100 ps every 153 ns [34]. At the Canadian Synchrotron Radiation Facility (CSRF) at the Synchrotron Radiation Center (SRC), University of Wisconsin-Madison and the Canadian Light Source (CLS) at Saskatoon, single bunch operation is required to provide sufficient time between bunches for decay measurements (300 ns and 570 ns for SRC and CLS, respectively) [35]. TRXEOL experiments are conducted in the same way as XEOL in the energy domain except that the PMT (photo-multiplier-tube) signal is time-gated with a selected time window [32–35]. Figure 8 shows the schematic for the TRXEOL measurements with multi-bunch (e.g., top-up at APS, 153 ns) and single bunch (e.g., CLS, 570 ns).


Figure 7 XEOL from porous silicon at the Si (~1840 eV) and Si oxide (~1847.5 eV) resonance; the corresponding TEY, PLY and the TEY of Si(100) are shown in the inset(for colours see online version)

Figure 8 Schematic for the time-structure of synchrotron light pulses. Left: electron buckets are filled evenly with a bunch length of ~3 cm (<100 ps) and the time separation between pulses is ∆t, ∆t = 153 ns at APS (top-up mode); right: single bunch operation in which only one of the buckets is filled. E.g., ∆t = 570 ns at the CLS (single bunch). The decay of the excited state via the optical mode is monitored within ∆t. The green arrow represents the synchrotron pulse and the red arrow represents the optical emission (for colours see online version)

3.2.3 XES (X-ray Emission Spectroscopy)

XES is a contemporary term for X-ray fluorescence, although it is often used to describe a shallow core hole being filled by an electron directly from the valence band, emitting a fluorescence photon in the process [36,37]. Thus, with the use of a monochromator to record the fluorescence X-rays, XES can probe the density of states of the valence band (via dipole transitions) that is often probed, alternatively by X-ray photoelectron spectroscopy. This technique is very desirable for the characterisation of nanostructures, which tend to charge up easily in XPS measurements. In addition, XES is more bulk sensitive. Together with the absorption spectrum, e.g., XANES, XES probes the occupied and the unoccupied electronic state on both side of the Fermi level, respectively. Experimentally however, XES demands a high brightness excitation source since for shallow cores of low Z elements, the fluorescence probability is significantly less than 0.1% [38]. With the advancement of the third generation synchrotrons and detection schemes, this technique is becoming very promising for routine applications. XES conducted at excitation energies in the vicinity of the absorption edge is in the

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regime of resonant inelastic X-ray scattering, which will provide additional information [39,40]. In this discussion however, we will limit ourselves to the three techniques and their applications to uniquely probe naonscale systems of interest. The schematic for both XEOL and XES has been shown in Figure 6. Within the single particle approximation, the fluorescence X-ray of the L3,2 shell of Si for example, probes the s character of the valence band.

4 Morphology, structure and electronic behaviour of selected nanosystems from synchrotron spectroscopy

In this section, we will present the analysis of several representative nano scale systems using a combination of conventional and synchrotron techniques. The systems of interest include Si and Ge nanowires and related systems, Au and other noble metal nanoparticles, II-V semiconductor nanoribbons, ZnS and CdS and related systems, ZnO nanowires and carbon based nanostructures such as nanowires, nanotubes and nanodiamond.

4.1 Porous silicon, Si and Ge nanowires and related systems

4.1.1 Porous silicon

The discovery of optical emission from porous silicon [1,2] around the turn of the 1990s and the discovery of Alq3, tris-(quinoline)aluminium, as an OLED (Organic Light Emitting Diode) material at about the same time [41] have contributed to the intense research and development in the area of nanotechnology relevant to optoelectronic applications and OLED technology that we are experiencing today. Porous silicon is produced by electrochemical etching of a silicon wafer in HF/ethanol solution, leaving behind a porous network of nodules and pillars of crystalline silicon of the dimension of nm. What is intriguing is that porous silicon emits light intensely at room temperature upon photo and electrical excitation and the energy of the optical emission is dependent upon the size of the Si nanostructure (the smaller the size the shorter the wavelength) while bulk silicon, an indirect band gap material (a radiative de-excitation requires the participation of a phonon) does not.

These exciting discoveries led to very intense research efforts in the decade to follow. It is now understood that light emission from porous silicon arises from an effect known as quantum confinement. In a particle-in-a-box picture, the separation between energy levels is inversely proportional to the dimension of the potential well. Thus quantum confinement results in the widening of the band gap and the PS discovery opened up the opportunity for band-gap engineering (wavelength tunable light emission). The use of synchrotron radiation played a significant role in identifying the origin of luminescence from porous silicon. Using the XEOL technique, we have shown that the origin of luminescence from porous silicon is due to quantum confinement and is site specific [30,42–45]. This can be illustrated in Figure 9 where X-ray excited luminescence from porous silicon prepared from a p-type Si (100) wafer is displayed.

From Figure 9, we see that the change of the branching ratio of the luminescence intensity at ~650 nm (nano Si) vs. ~480 nm (Si oxide) is apparent, indicating that the 650 nm luminescence is from porous silicon and the 480 nm from Si oxide [43,44].


HF treatment improves the luminescence intensity and causes a red shift as the result of the thinning of the oxide. The XEOL results for PS has recently been reviewed [46], we shall therefore focus below on a related system, silicon nanowires (SiNW).

Figure 9 Si L3,2-edge XANES of Si(100) and porous silicon (left panel) and luminescence from porous silicon excited at the Si L-edge resonance of the elemental silicon and above the silicon oxide L-edge before (mid panel) and after HF treatment (right panel)

4.1.2 Si nanowires

Recently developed techniques to fabricate silicon nanowires, a desirable configuration for nanodevice interconnects, have greatly enhanced the advanced characterisation of its electronic structure and optical properties [3–6]. All Si nanostructures have two features in common [2,30,43,47–49]:

a they all contain quantum confined elemental silicon structures of nm dimension

b they are often capped by silicon oxide.

As noted above, the light emitting properties of porous silicon are reasonably well studied [1,2]. The luminescence exhibits a range of size-dependent blue shifts relative to the indirect band gap energy of bulk silicon (1.1 eV) due to quantum confinement that results in the opening of the band gap [2]. In addition, the capping oxide layer is amorphous and contains sub-oxides and oxide, which also contribute to the luminescence [30,43]. It is expected that the light emitting properties of silicon nanowires are similar to that of other Si nanostructures. We summarise below our recent study of a SiNW specimen using the photon-in photon-out techniques: XEOL and XES as described above [49,50].

Figure 10 shows the XANES and XEOL results from SiNW prepared with techniques described in the literature [3]. The specimen has a nominal diameter of ~13 nm of which there is a silicon oxide outer-layer of ~3–4 nm as revealed by TEM. We can see from the left panel of Figure 10 that the luminescence spectra show three major broad bands at 460 nm (2.7 eV), 530 nm (2.34) and 630 nm (1.97 eV). The intensity and the branching ratio of the 2.7 eV peak reaches a maximum when the K-edge channel of silicon oxide turns on at 1847.5 eV. This is illustrated in inset (b) by the difference curve between the 1847.5 eV (SiO2) and the 1842 eV (Si) excited luminescence. We attribute the 2.7 eV peak to the luminescence from defects of the encapsulating oxide. This 2.7 eV luminescence is commonly observed in silica glass and has been attributed to the triplet-to-ground state luminescence from a SiO2 defect

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involving a bridging oxygen vacancy between two adjacent Si (≡ Si-Si ≡) [50–53]. It is interesting to note that luminescence excited at 1830 eV is also significant although this energy is below the Si K-edge. At this energy, photons are absorbed by the Si 2p, 2s and the valence electrons and the energy transfer to the optical channel is also efficient. Luminescence at similar wavelengths has also been observed in porous silicon and solution-grown SiNW [6,51–54]. It should be noted that below the SiO2 resonance and above the Si resonance, as shown in inset (a), elemental Si absorbs most of the photons. However, when the silicon oxide channel turns on, both the Si and SiO2 are competing for the incoming photons. The absorption at the SiO2 whiteline (the sharp peak at 1847.5 eV) will facilitate luminescence from the oxide chromophore if the process has a high quantum yield as is observed.

Figure 10 Left: XEOL of as-prepared SiNW excited with photon energies across the Si K-edge. Inset (a) and (b) show the Si K-edge TEY and the difference curve of XEOL excited at Si oxide and Si resonance. Right: Si K-edge XANES of SiNW recorded with TEY, FLY as well as zero order and wavelength selected luminescence

(a) (b)

The right panel of Figure 10 shows the Si K-edge XANES monitored with total (zero-order) and wavelength-selected luminescence (PLY). The TEY and X-ray fluorescence yield (FLY) are also shown. We see that in the surface sensitive TEY, the edge-jump ratio of silicon/oxide is ~1 : 1 compared to ~2.5 : 1 in FLY and zero-order PLY. The latter are good representations of the overall composition of the nanowire since these techniques are more bulk-sensitive. The most interesting features are seen in the wavelength-selected PLY where the 460 nm (2.7 eV) PLY exhibits the most intensity SiO2/Si ratio (1.3) whereas the 530 nm PLY the least (0.6). This observation strongly indicates that the luminescence at 460 nm (2.7 eV) is from silicon oxide while the 530 nm (2.33 eV) peak is primarily from small Si crystallites encapsulated in the SiNW. The emission is significantly blue-shifted relative to the indirect band gap of bulk silicon (1.1 eV) and that of porous silicon samples (600–700 nm or 2.07–1.77 eV) but is within range of the direct gap at the Γ point of the Si band structure and what has been observed


in many nanosilicon systems [1,51–54]. The crystallite size based on the quantum confinement model has to be significantly smaller than the diameter of the nanowire (a 2.33 eV gap would correspond to a ~2 nm crystallite), according to recent findings of experiment and theory [55–60]. Since the main core of the wire is larger than 2 nm, the Si based luminescence must arise from smaller nanocrystallites embedded throughout the wire, not the larger core. This also explains why in general luminescence from silicon nanowires is considerably weaker than that from porous silicon. The 630 nm luminescence is less sensitive to the excitation energy (SiO2/Si ratio ~0.8) and is attributed to interface defects between Si and silicon oxide (suboxide). This notion is also borne out in the increasing luminescence intensity going from excitation below to above the Si K-edge (~1842 eV) as shown in Figure 10. We shall return to this below.

The effect of surface chemistry on the optical properties of SiNW is shown in Figure 11 in which the XEOL from an aged SiNW specimen before and after HF treatment (partial oxide removal) is displayed. Before treatment, the XEOL is similar to those shown in Figure 10. After treatment, several spectral changes can be identified. First, as can be inferred from the SiO2 whiteline at 1847.5 eV, some oxide remains. Second, the overall luminescence intensity decreases after the HF treatment in which some samples were dissolved. Third and perhaps most surprising is the oxide luminescence at ~450 nm (blue-shifted by 10 nm from before HF) increases dramatically at 1847.5 eV. The 530 nm emission (red-shifted slightly) becomes less intense and the 630 nm peak is no longer noticeable. We attribute this observation to the loss of both Si nanocrystallites and silicon oxide due to HF treatment. However, HF must have removed a significant portion of the nano Si crystallites responsible for the 530 nm luminescence and produced, relatively, more oxide defects in the remaining SiNW. The oxide largely is responsible for the enhanced branching ratio and the site/excitation channel specificity observed in Figure 11.

We have investigated the valence electron density distribution of the SiNW using both XES and XPS (X-ray Photoelectron Spectroscopy) although XES was preferred over photoemission due to the charging problem of SiNW [50]. Figure 12 shows the XES excited above the Si L3,2 edge of the SiNW specimen before and after HF treatment together with that of a clean Si(100) and a SiO2 specimen (left). The XPS of the Valence Band (VB) obtained with monochromatised Al Kα for as–prepared PS, SiNW before and after HF, and Si(100) with the same energy scale (right) are also shown. The XES were excited with photon energy of 110 eV, which is above the Si L3,2 threshold of both elemental silicon and silicon oxide, and has an energy resolution of 150 meV. The broad band at the highest emission energy corresponds to the top of the valence band. The Si(100) XES shows a three-peak pattern characteristic of the sp band of bulk silicon below the Fermi level [36,58,59]. This is nearly identical to the valence band (VB) from photoemission of Si(100). The SiNW XES exhibits several interesting features. First, the as-prepared SiNW has a significant amount of oxide as can be seen in the region closest to the Fermi level. Second, there are states tailing towards the Fermi level, which diminishes upon HF treatment; we attribute them to interfacial states between Si and silicon oxide; this can be more readily observed in XES than in XPS since XES is more bulk sensitive. Finally, the Si XES features of crystalline Si emerge after HF treatment albeit some oxides remain on the surface. Similar features are also observed in the XPS VB. The slight difference in the relative intensity is due to difference in the cross-section. These observations confirm previous TEM results and the importance of the role of oxide layer in the growth of SiNW [29].

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Figure 11 Normalised XEOL of SiNW before (top) and after (bottom) HF treatment. The TEY XANES before and after HF treatment is shown in the inset. The spectra have been shifted vertically for clarity

Figure 12 XES excited at 110 eV photon energy (left) and XPS of the VB using Al Kα X-rays (right) for a series of Si nanostructures and Si (100)

It is also interesting to note that nano composites of SiNW containing alternate silicon and silicon oxide in a chain have also been produced by modulating the pressure of the carrier gas. One such an example is shown in Figure 13 [61].

From Figure 13, we see that the SiNW is composed of nodules of Si nanoparticles embedded in a silicon oxide chain with a pseudo-periodicity. HRTEM shows that the chain-like SiNW comprises the crystalline Si nanoparticles of an average diameter of 16 nm and ~20 nm-long interconnecting silicon oxide wires. In addition to the sphere, rectangle and triangle shapes of Si nanoparticles in chain-like SiNWs were observed.


Si K-edge XAFS and XEOL for these chains are shown in Figure 14 where one sees a significantly large Si oxide to elemental Si ratio in the edge jump of the chain-like SiNW (inset, left panel), consistent with the HRTEM observation and what is observed in Figure 11. XEOL (Figure 14 left panel) and PLY XAFS (Figure 14, right panel) show that strong luminescence from the chain-like SiNW arises primarily from the silicon oxide while quantum confined silicon luminescence is a minor component of the luminescence, in contrast to porous silicon luminescence where the Si nanonodules are smaller and contribute more significantly to the luminescence. It is evident from Figure 14(b) that the 460 nm PLY shows mainly SiO2 XANES features without any observable elemental Si resonance (~1840 eV) compared to the XANES recorded with zero order.

Figure 13 SEM (a) and HRTEM (b) and (c) of a SiNW prepared by modulating the flow of the carrier gas. In (c) Si NP embedded in an oxide chain is clearly visible

Figure 14 (a) XEOL from Si-SiO2 chain (top) and normal SiNW. The Si K-edge XANES is shown in the inset and (b) comparison of XANES recorded with TEY, FLY and PLY (zero order and 460 nm) yields

(a) (b)

Finally, we note that SiNW will be of great value as interconnect if we can increase the charge carrier mobility by doping. Our attempt to dope SiNW has met with some success [62]. The samples in this study were prepared using laser ablation with a mixture of SiO and phosphorus powder (1%). The TEM images shown in Figure 15 reveal the morphology of the nanowires.

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The doped SiNW was studied with XANES (TEY and FLY) at both the Si and P K-edge (Figure 16) at various stages of treatment with HF, which was carried out to remove the surface Si oxide. Let us compare the Si and P XANES. We see that the as-grown SiNW contains P, which was not removed by HF treatment (enhanced, in fact) indicating that P is contained in the SiNW core. The modulation of the XANES in energy is similar for both edges indicating that P is taking up a Si site in the nanowire.

Figure 15 SEM (left) and TEM of Phosphorus–doped nanowires

Figure 16 TEY and FLY spectra of K-edge NEXAFS after sequential HF etching: (a) TEY of Si K-edge; (b) FLY of Si K-edge; (c) TEY of P-edge and (d) FLY of P-edge

(a) (b)

(c) (d)


4.1.3 Ge nanowires

Compared to Si, renewed interest in Ge as a material of choice for future electronics is more recent and is due to its significantly higher electron and hole mobility relative to Si. These properties are desirable as electronic devices are scaled down to the sub-100 nm regime. Thus, Ge nanowires (GeNW) are expected to offer potential performance gains compared to Si-based planar and nanowire devices. Moreover, the exciton radius of Ge (24.3 nm) is larger than that of Si (4.9 nm), which consequently should lead to more prominent quantum confinement effects [63]. We used Au nanoparticles as the catalyst to prepare GeNW and benefited from the low eutectic temperature of Ge-Au alloy (360°C).The diameter and the location of the GeNWs on the substrate can be controlled by the size and location of the Au nanoparticles on the substrate. The scheme and product of this process have already been illustrated in Figure 1.

We have also grown GeNW on an Au–coated, patterned Si wafer. The direct self-assembly of GeNW onto the substrates with different properties or pattern with a catalyst will provide very useful information for future nano-devices applications. The electronic structure and local structure of GeNW, compared to Ge wafer, have also been investigated with XANES and EXAFS at the Ge K-edge. These techniques have been demonstrated to be powerful for the study of the electronic structure and optical properties of 1-D semiconductor nanomaterials [60,61,64].

Figure 17 shows the morphology of the GeNW with HRTEM image of the GeNW and the Au catalysts, which are found on the tip of the GeNWs [12,65]. The structure of the nanowire and the Au tip has also been confirmed with X-ray diffraction and Au L3-edge XAFS, respectively. It is interesting to note that these GeNW are weak light emitters in the visible upon X-ray excitation. More detailed analysis of XEOL from GeNW excited at the Ge L-edge and the O K-edge (GeNW is coated with an oxide layer) is in progress.

Figure 17 (a) HRTEM of a GeNW and its fast Fourier transform (FFT) image (inset) showing (111) diamond like crystalline planes and [111] growth direction and (b) HRTEM image of Au nanoparticle on the tip of a single GeNW showing fcc crystal structure

The growth of GeNW on a patterned surface is worth noting. Figure 18 shows the growth of GeNW on an etched Si surface with pits of ~5 µm size. The pits were loaded with an

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Au NP catalyst. SEM examination reveals that dense GeNW on the order of 10 nm only grows at and in the vicinity of the pits.

Figure 18 SEM of the etched substrate showing (a) µm size pits and (b) the morphology of the Ge nanowires

The Ge K-edge XAFS, XANES and EXAFS of GeNW and bulk Ge are shown in Figure 19(a) [12]. As in the case of Si noted above, the Ge K-edge XAFS of GeNW is very similar to that of bulk Ge except in details. The XANES for GeNW exhibits a sharp peak just above the edge attributable to a resonance associated with the surface oxide of the GeNW, analogous to SiNW. The EXAFS oscillations for GeNW has similar periodicity as that for bulk Ge albeit with noticeable broadening and reduction in amplitude as expected since the reduction of Ge-Ge coordination number on average and the degradation in long-range order in GeNW contribute to the reduction of the amplitude, A(k) (see equations (1) and (2)). This behaviour is clearly noticeable in the FT of the EXAFS and the filtered back-transform of the first shell radial distribution, shown in Figure 19(b) and inset. While the first shell Ge-Ge interatomic distance in GeNW appears to be similar to that of bulk Ge, there is a clear reduction in amplitude for GeNW, especially in the higher shells, i.e., that while the second and third shell of Ge-Ge (no phase correction) are clearly seen at ~3.6 and 4.3 Å in the R space for bulk Ge, the same cannot be stated for GeNW with confidence.

Figure 19 (a) Ge K-edge XAFS of GeNW and bulk Ge, the inset shows the XANES and (b) FT of the EXAFS, (k weighted, 3 section spline), the inset shows the filteredFT of the radial distribution of the first shell (for colours see online version)


4.2 Nanoparticles of metallic elements, surface vs. bulk effects in Au nanoparticles and related phenomena

In this section, we will review some recent work on the interplay of size and surface effects on the electronic structure of Au nanoparticle (NP). Since it is critical that size control must be obtained prior to systematic study, we will report below how this was carried out and confirmed. The important issues of these studies are

• How the structure and electronic properties of Au change in the nano regime as its size becomes smaller

• How important is the surface effect as the surface to bulk atom ratio increases

• What influence the capping molecules have on the electronic structure of the Au NP since nano Au particle is often stabilised by a variety of capping molecules from the very strong binding molecules such as thiol (-SH) to the weak binding functional groups such as hydroxyl (–OH) and amine (=NH).

Other issues such as the effect of the same capping molecule on different metal NP of the same size, or the quenching of luminescence from a light emitting substrate such as porous silicon and silicon nanowires, which are often used as a nanosubstrate to produce nanocatalyst of noble metals such as Rh and Pd, will also be addressed in this section.

4.2.1 Preparation, morphology and conventional analysis

The control of Au(III) ion and thiol concentration has produced nearly uniform Au NPs for systematic studies [15–17]. Figure 20 shows the TEM images of three groups of Au NPs of distinctly different size distribution prepared by colloidal methods with Au: thiol ratios of 6 : 1, 1 : 1 and 1 : 3. The TEM images in Figure 20(a)–(c) clearly show that the size of the NPs is distinctly different (distribution is ~±15% of the nominal size for 100 particles).

Figure 20 TEM images of thiol capped Au NP prepared with the concentration ratio of (a) 1 : 6; (b) 1 : 1 and (c) 1 : 3 and their distribution; the Au salt concentration is 0.01 M

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These specimens were analysed with XRD and UV-visible absorption. The XRD of the Au NPs is compared with that of bulk Au as shown Figure 21, which shows that the Au NPs exhibit diffraction patterns identical to fcc Au. A closed examination reveals that the 2θ shifts slightly to a larger value as the NP size decreases. This indicates a lattice contraction in the thiol capped NPs, in good agreement with recent observations [15,16]. The decrease in size is also seen in the increasing broadening of the XRD peaks. The diffraction peaks for the NPs are much broader than bulk Au and become broader still as the size of the NP decreases. The broadening can be interpreted in terms of the decreasing crystalline size (τ) in terms of the Scherrer equation [66]

/ cosK ττ λ β θ= (4)

where K is the shape factor with a value of 0.9, is the wavelength of the incident X-ray, and is given by (B2 – b2) in radian units where B is the width at half maximum of the diffraction peak and b is the diffraction peak width of corresponding bulk materials of which the contribution is mainly that of the X-ray beam height and the Darwin width of the crystal plane, and is the Bragg angle. The XRD results for this series of Au NP are in good accord with the TEM work.

The UV-visible absorption spectrum also shows the characteristic plasmon for metallic behaviour in the 4.0 nm NP but as the size decreases, the intensity of the plasmon peak decreases and vanishes in the 1.6 nm NP, an indication of metallic to insulator transition. We will return to this issue later.

Figure 21 XRD (left) and UV-visible absorption of a series of Au NPs shown in Figure 20

4.2.2 Analysis with X-ray spectroscopy

The XAFS (XANES and EXAFS) technique using synchrotron radiation contributes significantly to the characterisation of these Au NPs. Figure 22 shows the Au L3-edgeXAFS and the Fourier transform of the EXAFS for a series Au NPs of different size. Figure 23 shows the Au L3-edge XANES and the backtransform of the radial distribution function shown in Figure 22 with a window filtering out everything but the first shell.


Figure 22 Au L3-edge XAFS of Au NPs and corresponding FT of the EXAFS showing the appearance of the Au-S bond as the NP size decreases (for colours see online version)

Figure 23 Au L3-edge XANES and backtransform of the first shell radial distribution function(for colours see online version)

Figure 22 shows the characteristic fcc Au oscillations in the XAFS of the Au NPs although there is a noticeable degradation in the amplitude compared to bulk Au as the size of the NP decreases. As noted above, this behaviour is associated with the degradation of long-range order and reduction in coordination number. The Au L3-edgeXANES shows a noticeable increase in the whiteline intensity (first peak arising from 2p3/2 to 5d5/2,3/2 transition) as the size of the Au NP decreases. This observation immediately indicates that Au d charge at the Au NP is depleted on average. This depletion is due to the strong thiol–Au NP interaction. Analysis of the data using established procedures [67,68] shows that there is an increase of 11.2%, 9.0% and 7.2% in the d-hole count for the NPs of 1.6 nm, 2.4 nm and 4.0 nm respectively, relative to the Au metal where the d hole population (~0.40 from band structure calculation) above the Fermi level in Au arises from s-d hybridisation. This depletion of d charge in the Au d band leads to a partially unoccupied d band hence magnetic behaviour in these Au NPs. Our notion of thiol molecule induced holes in these Au NPs and the tuning of NP properties using the different influence of strong vs. weak capping molecules was recently demonstrated [69]. The results support this notion as illustrated in Figure 24 where the magnetic behaviour of strongly capped Au NP is shown.

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Figure 24 Magnetic measurement of Au NPs capped with weak interacting capping molecules such as amine (left) and strong interacting molecules such as thiol (right) [69]

We can see from Figure 24 that weakly interacting capping molecule (amine) does not induce magnetism while the strong interacting capping molecule (thiol) does as evident from the presence of the magnetic hysteresis loop. The size dependent behaviour of the Au 5d band in these Au NPs has also been studied [17]. It is well known that Au d-d interaction is very strong in metallic Au systems and that the Au 5d band width, apparent spin-orbit splitting ∆5d and the centroid of the d band are all very sensitive to the number of Au neighbouring atoms (coordination number). The d-band width W, and the apparent spin-orbit splitting, ∆5d are marked in Figure 25, which shows the XPS Au valence band spectrum of the Au NPs.

Figure 25 Valence band of Au NP. The vertical bars mark the bandwidth and the apparentspin-orbit splitting, ∆5d

The ∆5d can be expressed as the quadrature of the effect of band formation, ∆band, and the atomic spin-orbit splitting ∆so (1.5 eV) in Au atom,

2 2 1/ 25d band so( ) .∆ = ∆ + ∆ (5)


The band term decreases with the number of nearest Au neighbours. Thus the surface d band and apparent spin-orbit splitting are narrower than the bulk counter part [70]. It is established that when Au is diluted in alloy formation [70,71], both W and ∆bandbecome narrower as the coordination number of nearest Au neighbours decreases. We can see from Figure 25 that the top of the d band moves away from the Fermi level as the NP size decreases. This is accompanied by a bandwidth narrowing, a reduction in apparent spin-orbit splitting and a centroid shift (5.14 eV in bulk Au, 4.97 eV in 4.0 nm NP, 4.86 eV in 2.4 nm NP and 4.57 eV in 1.6 nm NP).

The Au NP-capping molecule interaction can also be investigated from the ligand’s perspective. The thiol-Au NP interaction as a function of size of the Au NP can be revealed by the S K-edge XANES shown in Figure 26, where the resonance resulting from Au-S interaction is apparent.

Figure 26 S K- and L3,2-edge XANES as a function of AuNP size in thiol (RSH)capped Au NPs

4.2.3 Unusual properties of molecular capped Au nanoparticles: coulomb blockade, luminescence and non-Fermi behaviour

Another interesting question concerning these Au NPs is whether or not at certain dimension, the NP becomes nonmetallic if the size of the NP continues to decrease. The metallic behaviour of materials is often gauged by its plasmon resonance, the collective motion of itinerant electrons, which can be measured experimentally using UV-visible absorption (as shown in Figure 19) or electron energy loss spectroscopy. The other gauge is the presence of a Fermi level, which can be measured with X-ray photoelectron spectroscopy (XPS) or a combination of XES and XANES although the latter presents some calibration challenges. In XPS, the Fermi level appears as a smoothed edge function at zero binding energy, and is often referred to as the Fermi edge, which results from the step function described by Fermi-Dirac statistics convoluted with the experimental resolution. Experimentally, it was often observed in metallic nano particle systems, which show plasmons, that the Fermi level was shifted away from zero to a higher binding and appeared skew. There has been some debate as to whether or not this shift implies a metal to insulator (semiconductor) transition or simply an experimental artifact arising from coulomb blockade (charging) by the capping molecule which separates the NP from the substrate such that the neutralisation of the XPS core hole by electron from the substrate is not fast enough to fully screen the core hole leading to a skew Fermi level [72–75]. This behaviour has been explained in terms of a time-dependent dynamic screening model [76].

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Still another issue concerning these Au NPs is that at what size, it becomes light emitting. There has been some very interesting reports on the size dependent photoluminescence of Au nano systems and at least in some cases, light emitting and metallic behaviour of these systems seem to be mutually exclusive [77,78]; i.e., that Au NPs that exhibit metallic behaviour do not exhibit photoluminescence whereas NPs’ that are light emitting do not exhibit metallic behaviour [79].

To address these issues, we have conducted an experiment with two Au NPs systems, one of which, an NP with a larger size, is distinctly metallic but not luminescent while the other, an NP with a smaller size, is luminescent but not metallic as judged by the presence/absence of plasmon and photoluminescence [79]. In the experiment, two sets of tiopronin-capped Au nanoparticles dispersed on a Si(100) substrate with a nominal size of ~3 nm (distribution in the range of 2–5 nm), and ~1.5 nm (all smaller than 2 nm) were used. They are henceforth denoted AuNP(L) and AuNP(S), for large and small NP, respectively. Tiopronin, CH3-CH(-SH)-CO-NH-CH2-COOH, is a water-soluble thiol. The tiopronin-capped Au NPs were prepared with established colloidal methods [14,18]. Upon capping, it allows the Au NP to be water soluble, a desirable property for chemical and biological sensor applications. Figure 27 shows the optical properties of these two systems.

Figure 27 UV-visible absorption of AuNP(L) (a) and AuNP(S) (b) and photoluminescence from AuNP(S) (c); AuNP(L) is not luminescent (not shown)

The conventional wisdom suggests that AuNP(L) is metallic and AuNP(S) is not. To test the other criterion, Fermi level, for metallic behaviour, we conducted an XPS measurement with a Kratos XPS instrument [80] that is equipped with a magnetic lenses system to minimise the differential charging of the system.

Figure 28 shows the Au 4f spectra of AuNP(S) and AuNP(L) as well as Au metal with neutralisation off (top left) and Au 4f of AuNP(S) with neutralisation on (bottom left). Right panel shows the valence band for AuNP(L) and AuNP(S) compared with Au metal. For AuNP(L) (top and mid panel), the Au d band width and spin-orbit splitting remain essentially unchanged with and without neutralisation while features at 14 eV and 11 eV were shifted by 1.3 eV. The bottom right shows the VB for AuNP(S) where all spectral features are shifted by 2.23 eV when neutralisation is on. The top of the valence band is now aligned at the zero binding energy (Fermi level).

Based on the above observation and the corresponding Au L3-edge threshold energy shift, a model has been proposed to explain the origin of the luminescence in small AuNPs. The schematic of the energy correlation is shown in Figure 29. The AuNP(S) spectra indicate that it is an insulator, thus there must exists a HOMO


(highest occupied molecular orbital) – LUMO (lowest unoccupied orbital) like pseudo gap. With neutralisation, the top of the valence band (HOMO) is aligned at the Fermi level (Figure 28). We propose a model with it we can estimate the gap using photoemission binding energy (relative to the HOMO) and X-ray absorption edge threshold energy (relative to LUMO).

Figure 28 Au 4f (left) and VB (right) of AuNP(L) and AuNP(S) with and without neutralisation. The top left panel also shows the Au 4f of Au metal and AuNP(L); the bottom left panel also shows the 4f of AuNP(S) without neutralisation shifted by –2.23 eV. The black curve on the right panel is the valence band of metallic Au (foil) (for colours see online version)

Figure 29 Schematic for the emergence of the pseudo gap as the size of the AuNP decreases below certain critical size. From left to right: Au (bulk), AuNP(L), AuNP(S)

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A schematic is given in Figure 29 where the energy separation of the gap can be estimated from the Au 4f binding energy shift and the Au L3-edge threshold shift between Au metal and AuNP(S).

3/ 23gap 2( ) ( )o pLE E S BE S∆ = − (6)

assuming 3/ 2 32 4 4( ) ( ) ( ) ( )p o f fLBE S E Au BE S BE Au− = −

3gap 4f .oLE E BE∆ = ∆ − ∆ (7)

Equation (7) assumes that the change in BE shift is the same for all core levels [81]. Note that for AuNP(S), the BE is referenced to the top of the valence band which lines up with the Fermi level of the substrate and the spectrometer when neutralisation is on. The Au 4f binding energy shift, ∆BE4f [AuNP(S)-Au] obtained from Figure 28 is –0.9 eV. The ∆Eo, L3-edge threshold shift [AuNP(S)-Au] (point of inflection of the rising edge, spectra not shown) yields a value of 0.8 ± 0.3 eV. Thus the pseudo gap we estimated for the AuNP(S) is ~1.7 (±0.3) eV. This value compares favourably with the luminescence maximum of 760 nm (1.63 eV) shown in Figure 27, indicating that the near energy gap emission is the origin of the luminescence.

4.2.4 Related phenomena: reductive preparation of Rh nano structureson nanosubstrates: porous silicon and silicon nanowires

Retuning to the hard template approach for the fabrication of nanoparticles, we use Rh to illustrate how synchrotron techniques can be used to characterise the nano particles (Rh) as well as the nano templates, porous silicon and silicon nanowires [64,82]. Figure 30 shows the XRD power diffraction of the Rh deposited on a HF refreshed porous silicon prepared with an n-type Si(100) wafer and the associated SEM images.

Figure 30 XRD for RhNP and Rh metal (left), and SEM of the substrate before (a) and (b)and after Rh deposition (c) and (d)

The XRD in Figure 30 shows a couple of interesting features that are common to nanostructures as was in the case of AuNPs. These features are

• broadening of the diffraction peaks

• reduction in intensity of the higher Miller index crystal planes.


Both features are signs of a nanostructure. The broadening has been noted earlier. When the (111) diffraction peak at 2 = 41.26º (d = 2.186 Å) in Figure 30 is used for the analysis (equation (5)), an average crystallite size of ~7 nm is obtained.

The surface of the as-prepared PS reveals macro-pore arrays consisting of square holes of tens to few hundreds of nanometres in size, and with edges lying along 110 directions. In addition, the pores penetrate directly down into the substrate as shown in Figure 30, right panel. Upon contact with the rhodium solution, the reaction immediately occurs on the surface. The surface of PS is now covered with a uniform film of Rh nanostructures. Figure 30, right panels (c) and (d), reveal that the film constitutes aggregates of Rh clusters with various sizes.

Figure 31 shows the Rh L3,2 edge XANES (TEY) of Rh metal and RhNP on PS. The XANES exhibit an intense whiteline (most intense peak at the edge arising from 2p3/2to 4d5/2,3/2 transitions) characteristic of unfilled states in the Rh d band just above the Fermi level. There is a noticeable difference in the oscillations in that features are broadened in RhNP. This indicates disorders in Rh nanostructures due to increasing surface to bulk atom ratio in small crystallites and associated strain resulting from the minimisation of the surface energy of a nanoparticle. Close examination shows that the L3–edge whiteline for RhNP/PS is more intense than that of the Rh metal. The ratio of the area under the L3 whiteline for RhNP to that of the Rh foil was about 1.6, which is surprisingly large. Thus there must be a significant contribution from the surface Rh atoms since the TEY at the Rh L3,2-edge is surface sensitive. We infer that there is some surface oxide (depletion in d-charge) which will contribute noticeably in the TEY whiteline at the Rh L3,2-edge.

Figure 31 Rh L3,2-edge XANES of Rh metal and RhNP on porous silicon

The Rh K-edge EXAFS (bulk sensitive due to hard X-ray and fluorescence detection) in Figure 32 clearly shows that RhNP/PS is metallic Rh. There is also a small positive threshold shift in RhNP/PS. This shift arises from the small size and surface effect of RhNP, in which a strong surface/interface interaction pulls electron charges away from the NP, resulting in a small blue shift. Similar behaviour is also observed for Rh deposited on Si nanowires (Figure 33) where amplitude reduction observed in the FT is also correlated with size. Co-deposition of Pd and Rh nanoparticles has also been successfully conducted as evident by the metallic features seen in Figure 33 (right panel).

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Figure 32 Rh K-edge XAFS and Fourier transform of the EXAFS of Rh metal and RhNP on porous silicon

Figure 33 EXAFS and FT for Rh metal and RhNP on SiNW prepared with different concentration (left) and metal L3,2-edge showing the co-deposition of Rh and Pd on SiNW (right)

We now examine the effect on the nano substrate. Figure 34 shows the XEOL from the porous silicon substrate before and after Rh deposition as well as the excitation energy-dependence of the luminescence.

Figure 34 (a) XEOL of ambient, HF-etched PS and RhNP/PS excited at 1841 eV and (b) XEOL from RhNP/PS excited with selected photon energy across the Si K-edge


From Figure 34(a), we see that XEOL is sensitive to the surface chemistry, and from Figure 34(b), we see that XEOL can be site, chemical and excitation channel specific as the selected photon energy scans across element Si and Si oxide thresholds. We can see that the branching ratio (intensity of an individual peak over total intensity) for the red emission (~790 nm) is the strongest at the Si edge resonance (1841 eV) whereas the branching ratio for the ~450 nm emission is the strongest at the SiO2 resonance (1848 eV), indicating that the Si and SiO2 are responsible for the luminescence at these wavelengths, respectively. These assignments are confirmed with the photoluminescence yield (PLY) of the XANES at both the Si K-edge and the Rh L3,2-edge (Figure 35).

Figure 35 PLY of RhNP/PS at (a) the Si K and (b) the Rh L3,2-edge

We can see from Figure 35 that PLY of zero order (all wavelengths), which is dominated by the red emission at the Si K-edge, shows an intense signal at the Si resonance but a dip at the SiO2 resonance and an inverted spectrum at the Rh L3,2-edge. This result indicates that the origin of luminescence is primarily Si and SiO2. Rh quenches the luminescence by competitive absorption.

4.3 Compound semiconductor nanostructures

In this section, we will review recent observations on the structure and electronic property correlation involving pseudo one-dimensional nanowire/ribbons of several compound semiconductors, CdS, ZnS, GaN and ZnO. These materials were all prepared with the thermal evaporation method with the use of a gold catalyst, the very thiol-capped Au nanostructures (~2 nm) we prepared using the soft template methods [17].

4.3.1 CdS and ZnS nanoribbons

One-dimensional semiconductor nanoribbons of hexagonal wurtzite sulfides (ZnS and CdS) have been prepared in bulk-quantity by a thermal evaporation technique using thiol-capped gold nanoparticles as catalyst [83–86].

Figure 36 (left panel) displays the TEM micrographs of the nanoribbons and the diffraction patterns from selected-areas of the same sample. It is interesting to note that both ZnS and CdS nanoribbons have a gold cap at one of the ends, indicating that these nanoribbons grew via the well-known vapour-liquid-solid (VLS) mechanism [7]. The strain resulting from the bending of the nanoribbon is evident. The thickness of the nanoribbons would not have exceeded several tens of nm as they are nearly electron transparent. The electron diffraction patterns for ZnS and CdS nanoribbon indicate

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that they have a hexagonal wurtzite structure. The high-resolution TEM images of a ZnS and a CdS nanoribbon (right panel) show that each nanoribbon is structurally uniform and single crystalline. For example, in the case of CdS nanoribbon, it grows along the (212) direction without stacking faults. Both high-resolution TEM and corresponding energy-dispersive X-ray (EDX) results confirm that these nanoribbons have a hexagonal wurtzite structure and they are indeed ZnS and CdS.

Figure 36 Left panel: TEM images of (a) ZnS and (b) CdS nanoribbon and the selected-area electron diffraction patterns of (c) ZnS and (d) CdS nanoribbon shown in (a) and (b), respectively. Right panel: HRTEM images of (a) ZnS and (b) CdS nanoribbon and corresponding EDX

Figure 37 shows the photoluminescence (PL) spectra of ZnS and CdS nanoribbons and their starting materials, ZnS and CdS powders (Alfa Aesar), excited with 300 nm and 320 nm photons (lab source) for ZnS and CdS respectively (top) and with synchrotron radiation (2450 eV and 2475 eV, bottom panel). Compared to powder sample, an additional peak at 343 nm and 516 nm for ZnS and CdS nanoribbon, respectively, is observed. The wavelength of these PL peaks is consistent with the band-gap energies for ZnS and CdS, respectively. XRD indicates that a phase transformation from zinc blende of the powder to wurtzite of the nanoribbon has occurred for ZnS, while CdS nanoribbon retains the wurtzite structure of the powder.

The XEOL spectra excited at photon energy above the S K-edge are shown in the bottom panel of Figure 37. The band-gap PL at 329 nm (3.8 eV) for ZnS nanoribbons and 528 nm (2.35 eV) for CdS nanoribbons is observed. In addition, intense peaks at 494 nm (2.5 eV) and 523 nm (2.4 eV) for ZnS nanoribbons, and at 445 nm (2.79 eV) for CdS, were also observed. The split intense peak in ZnS nanoribbon is due to partial oxidation (the maximum emission from ZnO nanostructure is at ~494 nm, see below) and 445 nm from CdS is attributed to states more accessible by X-rays. No XEOL could be detected for the powder. These results indicate that band-gap PL is enhanced markedly in nanoribbons.

Figure 38 shows the S K-edge XANES of the ZnS nanoribbon in both TEY and PLY (zero order) and its starting material, ZnS powders in TEY (no XEOL is observed). We see that there is an additional feature at ~2484 eV of the TEY of ZnS ribbon while this feature is absent in its PLY of ZnS nanoribbon and the TEY of ZnS powder. This feature has been attributed to surface S-O interactions and can be removed by sputtering [84].


Figure 37 UV (top) and synchrotron excited luminescence (bottom) for ZnS and CdS nanostructures and corresponding powder samples

Figure 38 S K-edge XANES in TEY and PLY for both nanoribbon and powder of ZnS (left)and CdS (right)

The zero order PLY XANES shows that for CdS nanoribbons the luminescence yield changes marginally across the S K-edge. This indicates that the luminescence from the nanoribbon is more likely originated from Cd than S [87,88].

4.3.2 GaN nanowires

The preparative and characterisation methods described above have been extended to the study of GaN nanowires [89]. GaN is a semiconductor with a high melting point,

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large carrier mobility, high electrical breakdown field, and a wide band gap (3.4 eV), and is of intense interest for ultraviolet (UV) or blue light emitters, and high-performance, high-power electronic devices [90,91] applications. This system displays two distinctly different zigzag morphologies, I and II, which can be inter-converted as shown in Figures 39 and 40, respectively.

Figure 39 Top: TEM and HRTEM of the type I zigzag-shaped 1D GaN nanostructures shownin (A) and (B), respectively, in (C) is the selected-area electron diffraction pattern; electron diffraction pattern was indexed to the reflection of wurtzite GaN [2 1 10]

Figure 40 Left: (A) SEM and (B) HRTEM images of a zigzag 1D GaN nanostructure (type II). The angle formed by the outline of the TEM image is 117°. The nanostructure is also a single crystal and defect free, and grows along [0001]. The angle between the two edges is 124°, which is equal to the angle between two{10 11} planes, {10 11} and {10 11}, for example. Right: (A) Transformations or junctions between two types of zigzag-shaped 1D nanostructures and (B) Mirror-symmetric zigzag-shaped nanostructure with a ‘defect’ (M is the mirror plane)


The GaN nano zigzag structures are also light emitting albeit with weak intensity. Figure 41 shows the XEOL of GaN nanostructures excited at 1000 eV. We see that the emission is dominated by a blue light ~400 nm with some red in the region of 650–800 nm. The blue emission is consistent with band gap emission from GaN. Some of the emission intensity at 756 nm is due to second order light.

Figure 41 XEOL from GaN zigzag nanostructures showing blue and red luminescence

4.3.3 ZnS nanoribbons: XEOL, TRXEOL and phase specificity

The ZnS nanostructure has been subjected to a temperature dependent study in which HRTEM results were used to help identify the origin of the luminescence. Figure 42 is the HRTEM image of the ZnS nanoribbons showing that the majority of the ZnS nanoribbon is hexagonal wurtzite, but there is an area close to the Au cap where cubic sphalerite can be identified [86,92].

Figure 42 (a) High-resolution TEM image of the boxed region of the ZnS nanowire in the region near the Au particle (b) showing sphalerite, and wurtzite structures as shown by the selected area diffraction patterns in (c) and (d), respectively

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The XEOL from ZnS nanoribbons (Figure 43) exhibits three emission peaks at 338 nm, 430 nm and 520 nm. Both the 338 nm and the 430 nm emission are enhanced significantly at low temperature. This result indicates that they are coupled effectively to the phonon and are quenched effectively at room temperature. The Zn L-edge XANES monitored with the optical yield at these wavelengths shows different results: the 338 nm PLY resembles that of the wurtzite structure while the 430 nm and 520 nm PLY resembles that of the sphalerite [92]. The 338 nm peak is associated the dominant wurtzite structure while the 430 nm and 520 nm are associated with vacancy and Au impurity in the sphalerite structures, respectively.

Figure 43 (a) XEOL from ZnS nanoribbons excited at 1000 eV at 20 and 280 K showing emissions at 338 nm, 430 nm and 520 nm. The inset shows the near-band gap region. (b) Total fluorescence yield (TFY) XANES of ZnS nanoribbon compared with that of ZnS (wurtzite). (c) and (d) show the wavelength-selected PLY and the XANES of ZnS nanoribbon at the Zn L3,2-edge compared with those of sphalerite and wurtzite(for colours see online version)

4.3.4 ZnO nanowires, nanoribbons and nanorods

ZnO is a promising candidate for optoelectronics and piezoelectricity effects due to a wide band gap (3.37 eV), a large exciton binding energy (60 meV) at room temperature and its non-centrosymmetry in the wurtzite structure [93,94]. A large variety of ZnO nanostructures, such as nanowires (nanorods) [95], nanotube [96], nanobelts [97], nano combs [98], nanosprings [99], nanopropellers [100], and tetrapods-like nanostructures [101,102], has been fabricated using a variety of techniques. Here we report a simple approach to obtain different ZnO nanostructures such as single crystal nanowires, nano-needles, nanoflowers and tubular whiskers via controlled oxidation of metallic Zn powder without the use of any metal catalyst at a modestly low temperature of ~550°C. We show that specific ZnO nanostuctures can be obtained at a specific temperature zone in the furnace by controlling the temperature and the partial pressure of oxygen and Zn. We show below that XEOL from these ZnO nanostuctures exhibit noticeable morphology-dependent luminescence associated with crystallinity, defect concentration and surface roughness.

The ZnO nanostructures discussed here were synthesised on a Si(100) substrate in a horizontal tube furnace. The morphologies and microstructures of the products at different temperature zones are shown in Figures 44–46.


Figure 44 Scanning electron microscopy (SEM) images of different ZnO nanostructrues obtained at different positions on the Si substrates. (a) ZnO nanowires; (b) ZnO particles;(c) ZnO nano-needles and (d) ZnO nano-flowers

Figure 45 (a) TEM image of several ZnO nanowires, which have similar diameter of around15 nm and (b) HRTEM images of the body of a single ZnO nanowire

Figure 46 (a) TEM image of a two-segment nano-needle and its corresponding electron diffraction pattern (inset) and (b) HRTEM images of the tip

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From Figure 44, we see very different morphology from different temperature zones [93]. From Figures 45 and 46, we see that there is a marked difference in the crystallinity of the nanostructures. The nanowire is bent and full of defects while the nano needle is a nearly perfect crystal. These structural differences are borne out in the luminescence spectra (Figure 47) where there is a dramatic variation in the branching ratio (individual band intensity over total emission) of the band gap (~380 nm) and green emission (485 nm and 520 nm), i.e., that the nano needle exhibits an intense band gap emission with the green emission only barely visible while the nanowire exhibits a opposite behaviour with an intense green band and a weak near bad gap emission. This observation confirms a general notion of intrinsic luminescence; i.e., that the intensity of intrinsic luminescence is correlated to crystallinity.

Figure 47 (a) XEOL spectra from ZnO nanowires, nano-needles, nanoparticles and bulkZnO pure powder; the area under the entire spectrum has been normalised to unity. (b) The branching ratio of the band-gap emission and the green light (defect) as a function of decreasing crystallinity (left to right) (for colours see online version)

4.4 Carbon based nanostructures and related studies

The final illustration deals with the study of carbon-based nano materials. The discovery of C60 [103] and carbon nanotubes [104] has created a new field for research. We will present several examples below to show that XANES at the carbon K-edge can be very useful in probing the structure, electronic properties and texture of carbon-based materials taking advantage of the intense, tunable and polarised soft X-rays from synchrotron radiation.

Figure 48 shows the structure of graphite and the C K-edge XANES as a function of polarisation (angle of incidence). Synchrotron light from a bending magnet or a linear insertion device is nearly 100% linearly polarised in the plane of the electron orbit. The intensity, I of the transition from a core |ið to a bound state or a quasi-bound state (multiple scattering state supported by a potential barrier), | f ð, such as a π* or a σ*resonance, respectively, is described by

2 2cosI i r f i fε ϑ∝ ⋅ ∝ (8)

where ε and r are the polarisation (electric field) vector of the photon and the electric dipole vector of the system (electronic state of interest), respectively. The angle θ is the angle between the ε vector and the molecular orbital axis. In this context, the molecular orbital axis, or r, is often defined as the axis along which, the amplitude of the


wavefunction is a maximum. This is to be distinguished from the angle δ in Figure 48, which is the angle between the ε vector and the surface normal [105,106].

Figure 48 Structure of graphite along the c axis and normal incidence of the X-ray photon (left) and angle-dependent XANES of graphite at the C K-edge showing the polarisation dependence of the π* and σ* resonance (right); δ is the angle between the polarisation vector and the surface normal. 90° means normal incidence

Graphite has a set of coplanar, delocalised π* unoccupied electronic states of largely p character (sp2 hybridisation) normal to the plane and a set of in plane σ* unoccupied states. Note that the orbital axis for π* is normal to the plane of the graphite sheets and that for the σ* orbital is in-plane. Thus they are orthogonal. When a sample is mounted vertically as is typically done, the ε vector of the photon is in the plane of the sample (normal to π* orbital) at normal incidence and nearly normal to the surface at glancing incidence (nearly normal to the σ* orbital). Thus the δ in Figure 48 is the same as θ for the 1s – π* transition but δ = 90 – θ for the 1s – σ* transition.

From equation (8), we can see that the intensity of the transition is at maximum if the angle between the ε vector and the orbital axis is parallel (θ = 0°). This is clearly seen in Figure 48 where the π* resonance increases as the angle of incidence becomes more glancing at the expense of the σ* resonance intensity. Since the carbon nanotube can be visualised as resulting from the rolling graphene sheet, it will exhibit texture effects depending on the orientation of the rolling. Angle-dependent C K-edge XANES can probe these electronic structure and texture effects.

4.4.1 Carbon nanowires (CNW) and multi-wall carbon nanotube (MWNT)

One interesting property of the carbon nanotube (CNT) is that it can be prepared from amorphous carbon nanowires (CNW). Figure 49 shows the high-resolution TEM image of the CNW and CNW-converted multi-wall carbon nanotube (MWNT) used in the XANES measurements. It has been shown that CNWs can be prepared by thermal evaporation; the CNW thus prepared can be subsequently converted to MWNT by heating [107,108]. From Figure 49(a), we see that there are no lattice fringes inside the CNW unlike MWNT in Figure 49(b) where the nearly parallel dark lines (separated by 0.34 nm) in the MWNT are clearly visible; these are the graphene basal planes. The inner and the outer diameters of a typical nanotube are 3 nm and 20 nm, respectively. Discontinuity and defects in the nanotube are also apparent.

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Figure 49 TEM images of carbon nanowire CNW (left) and CNW-converted multiwall nanotube, MWNT (right)

The C K-edge NEXAFS of CNW and the CNW-converted MWNT in TEY and FLY are given in Figure 50(a) together with the TEY of HOPG (highly ordered pyrolytic graphite) in normal incidence (to the substrate onto which the samples were deposited). It should be noted that TEY and FLY at the carbon K-edge has probing depth on the order of nm and 102 nm, and are hence surface and bulk sensitive, respectively. Resonance at ~285 eV and ~286 eV in the TEY of both CNW and MWNT are attributed to C 1s to π* transitions which are also observed in the FLY albeit somewhat blurry due to self-absorption. These features immediately indicate the presence of unsaturated carbon-carbon interactions (sp2 bonding) [25,105,106], delocalisation and the presence of chemically non-equivalent carbon. The Β* transitions in CNW-converted MWNT are more pronounced than that of CNW. The most interesting observation is in the region of 287–290 eV where CNW exhibits intense features in the TEY, which are suppressed in the FLY. A similar feature is also detected at ~289.5 in the TEY of MWNT but not in the FLY. These are Φ* resonance characteristics of C-H bonds in amorphous carbon and are attributable to C-H bonding in both CNW and CNW-converted MWNT (Figure 49). The low energy shoulder in CNW contains some sp3 contribution that vanishes in MWNT. Except for the C-H Φ* resonance in the TEY from the residues, the NEXAFS of the CNW-converted MWNT are in very good accord with previous MWNT results [109,110]. This observation strongly indicates that CNW, despite its generally amorphous nature revealed by HRTEM, already contains building block seeds (graphitic carbon) inside. These units ultimately lead to the formation of MWNT upon annealing.

The presence of local order and texture in the CNW is observed in the angular dependence of the XANES of the CNW. Since the synchrotron light is linearly polarised horizontally, the intensity of the π* transition is sensitive to the orientation of the π*orbital with respect to the polarisation vector as noted above. Thus if the π* orbitals in the nanowires are partially oriented, a rotation of the specimen with respect to the incoming photon will show a measurable angular dependence. A C-C or C-H Φ* orbital orthogonal to the π* on the other hand will show an opposite trend. Figure 50(b) shows the angular dependence of the XANES of CNW in both TEY and FLY: the π* resonance intensity decreases from normal to glancing incidence while the C-H σ* exhibits an opposite trend. The intense C-H σ* resonance observed at glancing angle indicates the presence of C-H bond in the graphitic-like basal plane. This observation indicates partial orientation (texture) of the precursor inside the CNW such that the Β* orbitals are partially oriented and orthogonal to the C-H Φ*. Hence the MWNT building block within CNW is likely to have the structure somewhere between a phenyl group and a graphene-like basal plane (sp2 bonding).


Figure 50 (a) C K-edge XANES of CNW, MWNT and HOPG and (b) angle dependent XANESof CNW

(a) (b)

4.4.2 Nanodiamond and BN films

Another technologically important allotrope of carbon is diamond. Recent advances in diamond synthetic techniques both in bulk and in nano dimensions have generated considerable excitement [111,112]. Here we highlight the C-K-edge XANES of nanodiamond prepared by the HFCVD method with of a mixture of methane and hydrogen.

The specimens were prepared by a custom-designed, double bias assisted HFCVD (Hot Filament Chemical Vapour Deposition) equipment as reported previously [113,114]. TEM of the nanodiamond film clearly shows that nanodiamond particles with an average diameter of 10 nm are embedded in an amorphous carbon matrix. The surrounding amorphous carbon matrix plays an important role in nanodiamond growth, since nanodiamond films produced by the conventional HFCVD tend to become micro-scale [115]. TEM data analysis reveals {111} crystalline planes with a lattice spacing of 0.21 nm, consistent with the standard {111} plane of diamond.

The C K-edge XANES of a nanodiamond film recorded in TEY is given in Figure 51(a); spectra of CVD diamond, diamond-like carbon (DLC) and highly ordered pyrolytic graphite (HOPG) are also shown. Similar to previous results of CVD diamond [116], the spectrum of the nanodiamond film exhibits a sharp spike at 289 eV (excitonic transition), and a dip at 303 eV, both characteristics of diamond [117]. Close observation (inspecting the first derivative) reveals that the nanodiamond film exhibits a slightly broadened exciton transition with a 0.25 eV blue shift (Figure 50(b)). The blue shift indicates a reduced exciton radius or the widening of the band gap, a characteristic behaviour of quantum confinement in nanosystems of semiconductors [30,116]. The broadening is associated with the distribution of the size of the nanocrystallites. There is also a weak peak at 285.5 eV in the spectrum of the nanodiamond film that is assigned to the π* feature (sp2) due to the unsaturated carbon in the film. A more intense 285.5 eV peak was observed for the HOPG and the DLC samples as expected. The above observation indicates that the nanodiamond consists of predominately nanodiamond

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crystallites with a small amount of sp2 type carbon, in agreement with the TEM results. In addition, the resonance beyond the exciton peak is progressively farther apart in the nanodiamond XANES relative to that of the CVD, indicating a possibly contracted lattice in the nanodiamond.

Figure 51 Left: C K-edge XANES of diamond and nanodiamond compared with DLC (diamond like carbon) and HOPG. Right: Comparison of micro and nano diamond XANES in the vicinity of the edge

The nanodiamond thin film has also been studied by an angular dependent technique taking advantage of the linear polarisation of the synchrotron light. This technique can detect the texture of the film [25,105,106]. Figure 52(a) shows the angular dependent spectra of the nanodiamond film, recorded in the TEY. The angle marked in the figure is the angle between the incident light and the surface of sample. For example, in the case of HOPG, the wave-vector of the light is parallel to the surface at normal incidence (see Figure 48). This orientation enhances the σ* transition and inhibits the π* transition. It is clear that there is no noticeable angular dependence, suggesting that the film is polycrystalline with no preferred orientation. This is not surprising given the high local symmetry of c in diamond. It is also interesting to compare the surface sensitive TEY spectra with the bulk sensitive FLY spectra shown in Figure 52(b). It is obvious that the surface of the sample is more diamond-like than the bulk. This observation suggests that a significant amount of the amorphous carbon matrix remains intact during the production of the film with only the surface amorphous carbon converted to nanodiamond crystallites.

Figure 52 Angle dependent C K-edge XANES of nanodiamond: (a) TEY and (b) FLY


The XANES results reported here are in good agreement with the previous studies [111,116,118]. However, the starting material in the previous work was C60 and it was believed that hydrogen-poor plasma was a key. This method has been regarded as the standard method to synthesise nanodiamond thin film. Our results show that in addition to the hydrogen-poor plasma technique, the hydrogen-rich plasmas technique is also effective.

Finally, the blue-shifted excition observed in our experiment is in good accord with those reported earlier [116,117] except that our observed shift of 0.25 eV for a size of ~10 nm is smaller than that predicted by the previous observation which would suggest a value of >0.4 eV for this size. The discrepancy may result from inhomogeneity, different surface effect and morphology of the crystallites.

Another diamond-like material of interest is cubic boron nitride (c-BN), which is the second hardest material next to diamond. These two systems have been widely studied in the past decade due to their extraordinarily desirable physical and chemical properties. Cubic BN has the properties comparable and in some aspects even superior to that of diamond. Many HRTEM works show that c-BN grows in a sequence of amorphous (a-BN), turbostratic BN (t-BN), and then c-BN layers (both a-BN and t-BN are sp2- hybridised but in different structures with the latter more graphite-like), A simplified schematic is shown in Figure 53 where the t-BN is found with TEM studies to be well oriented with its (0002) basal planes perpendicular to the Si substrate. This system, which has a surface layer of largely cubic structure and a texture layer underneath (sp2 bonding, graphite-like) of which the XANES is sensitive to polarisation, presents a unique opportunity to illustrate the power of XANES to reveal the texture of underlayer [119].

As noted above, the fluorescent X-ray is bulk sensitive because of its relatively long attenuation length (e.g., 100–400 nm in c-BN near B K-edge and a little longer near N K-edge) [120]. If the thickness of a film is less than the penetration length of the incident and the fluorescent X-ray, the information from the substrate-film interface can be obtained by FLY. Up to now, literatures on the XANES of c-BN films is relatively lacking [121]. We have conducted an angular dependent B and N K-edge XANES study of c-BN thin films with various thicknesses on Si. The XANES are shown in Figure 54(a) and (b). Other related systems can be found in the literature [122].

Figure 53 Schematic of a c-BN film where the surface is primarily cubic with a turbostratic layer underneath and a amorphous layer at the BN Si interface

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Figure 54 Angle dependent B K-edge XANES of a BN film in (a) TEY and (b) FLY. The dominant species are marked. The corresponding peaks are labelled from 1 to 4(for colours see online version)

Two interesting features from the Figure 54 deserve some attention. First, the boron K-edge 1s – π* resonance (analogous to π* resonance in graphite), is very weak in the TEY, indicating that the surface layer is primarily c-BN. Second, the π* resonance in FLY, which is sensitive to the underlayer, is very intense, and exhibits angular dependence. This observation confirms that the underlayer has considerable unsaturation (sp2 bonding) and exhibits preferred orientation. Detailed analysis has been carried out to quantify the results. The perhaps most intriguing observation is that the underlayer and its texture can be probed with FLY and in this case, it is significantly different from the surface layer. This result also confirms the notion of c-BN growth as shown in Figure 53.

5 Summary and prospects

We have presented several very powerful characterisation tools based on synchrotron radiation spectroscopy and their applications in unravelling secrets from nanostructures. We have illustrated with examples that XAFS, XANES, EXAFS, XES, XPS and XEOL can be used to elucidate the structure and electronic properties of representative metal nanoparticle, Au and Rh NPs and semiconductor nanowires, Si, Ge, CdS, ZnS, GaN, ZnO and carbon nanotubes. We show that many properties that cannot be easily probed with conventional techniques can be investigated readily with synchrotron spectroscopy techniques. While these techniques will play an increasingly significant role in nanostructure characterisation for years to come, with the advancement in synchrotron technology, timing measurements will emerge to become another novel synchrotron tool in the characterisation of the optical properties of nanostructures, transient structures and dynamics of energy transfer in the excited state. We will end with a recent TRXEOL study of SnO2 nanoribbon at the CLS using a single bunch (Figure 55). The results demonstrate that timing measurements can reveal hidden optical emission. In Figure 55, peak C, recorded with a fast time window (0–10 ns) is not readily observable by XEOL in the energy domain (un-gated) in which only peaks A and B are apparent [123]. TRXEOL as well as its applications in the study of nanomaterials was recently reported elsewhere [124].


Figure 55 TRXEOL from SnO2 nanoribbons showing a hidden emission at ~450 nm with a fast time gate (0–10 ns). This peak (C) is not readily seen in the ungated (0–570 ns) or slow-window gated XEOL (for colours see online version)

Acknowledgements

Many of my former and present students, postdoctoral fellows and collaborators have contributed to the work described here. I am particularly in debt to Ian Coulthard, Steve Naftel, Detong Jiang, R. Sammynaiken, Peng Zhang, Grace Kim, Terry Tang, Jeff Sun, Xingtai Zhou and Yun Mui Yiu of UWO. The staff at the Canadian Synchrotron Radiation Facility (CSRF) over the years, Kim Tan, Brian Yates, Yonfeng Hu, Astrid Jurgensen, Franziskus Heigl, Canadian Light Source staff Tom Regier, Robert Blyth and PNC-CAT staff Robert Gordon have also played a significant role. I am also grateful to many of my collaborators, particularly Professors S.T. Lee and N.B. Wong of the City University of Hong Kong and Drs. Richard Rosenberg and Gopal Shenoy of Argonne National Laboratory and Professor Alex Moewes of the Physics Department, University of Saskatchewan for their collaboration. The research conducted at the University of Western Ontario was supported by NSERC, CFI, CRC of Canada and OCMR, MMO and OIT of Ontario. The Canadian Synchrotron Radiation Facility at the Synchrotron Radiation Center (SRC), University of Wisconsin-Madison was supported by NSERC and NRC. Work at the Canadian Light Source was supported by NSERC, NRC, CIHR and the University of Saskatchewan. The PNC-CAT of APS was supported in part by an NSERC MFA grant. SRC and APS were supported by the US National Science Foundation and the Department of Energy, Office of Basic Sciences, respectively.

References and Notes 1 Cullis, A.G., Canham, L.T. and Calcott, P.D.J. (1997) ‘The structural and luminescence

properties of porous silicon’, J. Appl. Phys., Vol. 82, No. 3, pp.909–965. 2 Canham, L.T. (1990) ‘Silicon quantum wire array fabrication by electrochemical and chemical

dissolution of wafers’, Appl. Phys. Lett., Vol. 57, No. 10, pp.1046–1048. 3 Zhang, Y.F., Tang, Y.H., Wang, N., Yu, D.P., Lee, C.S., Bello, I. and Lee, S.T. (1998),

‘Silicon nanowires prepared by laser ablation at high temperature’, Appl. Phys. Lett., Vol. 72, No. 15, pp.1835–1837.

1240 T.K. Sham

4 Morales, A. and Lieber, C.M. (1998) ‘A laser ablation method for the synthesis of crystalline semiconductor nanowires’, Science, Vol. 279, pp.208–211.

5 Wang, N., Tang, Y.H., Zhang, Y.F., Yu, D.P., Lee, C.S., Bello, I. and Lee, S.T. (1998), ‘Transmission electron microscopy evidence of the defect structure in Si nanowires synthesized by laser ablation’, Chem. Phys. Lett., Vol. 283, Nos. 5–6, pp.368–372.

6 Lee, S.T., Wang, N., Zhang, Y.F. and Tang, Y.H. (1999) ‘Oxide-assisted semiconductor nanowire growth’, MRS Bulletin, Vol. 24, No. 8, pp.836–842.

7 Wagner, R.S. and Ellis, W.C. (1964) ‘Vapor-Liquid-Solid mechanism of single crystal growth’, Appl. Phys. Lett., Vol. 4, No. 5, pp.89–90.

8 Hillert, M. (1998) Phase Equilibra, Phase Diagram, Phase Transformation, Cambridge University Press, Cambridge.

9 Nguyen, P., Ng, H.T. and Meyyappan, M. (2005) ‘Growth of individual vertical germanium nanowires’, Adv. Mater., Vol. 17, No. 5, pp.549–553.

10 Mathur, S., Shen, H., Sivakov, V. and Werner, U. (2004) ‘Germanium nanowires and core-shell nanostructures by chemical vapor deposition of [Ge(C5H5)2]’, Chem. Mater., Vol. 16, No. 12, pp.2449–2456.

11 Korgel, B.A. and Hanrath, T. (2003) ‘Supercritical fluid–liquid–solid (SFLS) synthesis of Si and Ge nanowires seeded by colloidal metal nanocrystals’, Adv. Mater., Vol. 15, No. 5, pp.437–440.

12 Sun, X.H., Didychuk, C., Sham, T.K. and Wong, N.B. (2006) ‘Germanium nanowires: synthesis, morphology and local structure studies’, Nanotechnology, Vol. 17, pp.1–6.

13 Coulthard, I., Zhu, Y-J., Degan, S. and Sham, T.K. (1998) ‘Reductive deposition of Au on porous silicon: preparation, morphology and electronic structures’, Can. J. Chem., Vol. 76, No. 11, pp.1707–1716.

14 Zhang, P. and Sham, T.K. (2003) ‘Fabrication of thiol-capped Pd nanoparticles: an electrochemical method’, Appl. Phys. Lett., Vol. 82, No. 8, pp.1778–1780.

15 Zhang, P., Kim, P.S. and Sham T.K. (2002) ‘Electrochemical route for the fabrication of alkanethiolate-capped gold’, Appl. Phys. Lett., Vol. 82, No. 9, pp.1470–1472.

16 Zhang, P. and Sham, T.K. (2001) ‘Tuning the electronic behavior of Au nanoparticles with capping molecules’, Appl. Phys. Lett., Vol. 81, No. 4, pp.736–738.

17 Zhang, P. and Sham, T.K. (2003) ‘X-ray studies of the structure and electronic behavior of alkanethiolate-capped gold nanoparticles: the interplay of size and surface effect’, Phys. Rev. Lett., Vol. 90, No. 24, pp.245502(1–4).

18 Zhang, P., Zhou, X., Tang, Y. and Sham, T.K. (2005) ‘Organosulfur functionalised Au, Pd and Au-Pd nanoparticles on 1D silicon nanowire substrates: preparation and XFAS studies’, Langmuir, Vol. 21, No. 18, pp.8502–8508.

19 Winick, H. (Ed.) (1994) Synchrotron Radiation Sources, A Primer, World Scientific, Singapore.

20 Attwood, D. (2000) Soft X-ray and Extreme Ultraviolet Radiation, Principle and Applications,Cambridge University Press, Cambridge.

21 Sayers, D.E., Stern, E.A. and Lytle, F.W. (1971) ‘New technique for investigating noncrystalline structures: Fourier analysis of the extended X-ray absorption fine structure’, Phys. Rev. Lett., Vol. 27, No. 18, pp.1204–1207.

22 Lee, P.A., Citrin, P.H., Eisenberger, P. and Kincaid, B.M. (1981) ‘Extended X-ray absorption fine structures- its strength and limitations as structure tool’, Rev. Mod. Phys., Vol. 53, No. 4, pp.769–806.

23 Koningsberger, D.C. and Prins, R. (1988) X-ray Absorption Principles, Applications, Techniques of EXAFS, SEXAFS and XANES, John Wiley and Sons, New York.

24 Bianconi, A. (1988) ‘XANES spectroscopy’, in Koningsberger, D.C. and Prins, R. (Eds.): X-ray Absorption Principles, Applications, Techniques of EXAFS, SEXAFS and XANES,John Wiley and Sons, New York, pp.573–662.


25 Stöhr, J. (1988) NEXAFS Spectroscopy, Springer-Verlag, New York. 26 Rehr, J.J. and Albert, R.C. (2000) ‘Theoretical approaches to x-ray absorption fines structure’,

Rev. Mod. Phys., Vol. 72, No. 3, pp.621–654. 27 Zhang, Y.F., Liao, L.S., Chan, W.H., Lee, S.T., Sammynaiken, R. and Sham, T.K. (2000)

‘Electronic structure of silicon nanowires: a photoemission and X-ray study’, Phys. Rev. B,Vol. 61, No. 12, pp.8298–8305.

28 Sham, T.K. (1986) ‘Application of X-ray absorption spectroscopy to the studies of structure and bonding of metal complexes in solution’, Acc. Chem. Res., Vol. 19, No. 4, pp.99–104.

29 Rogalev, A. and Goulon, J. (2002) ‘X-ray excited optical luminescence’, in Sham, T.K. (Ed.): Chemical Applications of Synchrotron Radiation, Part II: X-ray Applications, World Scientific Publishing Co., Singapore, Vol. 12B, pp.707–760.

30 Sham, T.K., Jiang, D.T., Coulthard, I., Lorimer, J.W., Feng, X.H., Tan, K.H., Fraigo, S.P., Rosenberg, R.A., Houghton, D.C. and Bryskiewicz, B. (1993) ‘Origin of luminescence from porous deduced by synchrotron-light-induced luminescence’, Nature, Vol. 363, pp.331–334.

31 Hu, Y.F., Tan, K.H., Kim, P.S., Naftel, S.J., Sham, T.K., Coulthard, I. and Yates, B.W. (2002) ‘Soft X-ray excited optical luminescence: some recent applications’, Rev. Sci. Instru., Vol. 73, No. 3, pp.1379–1381.

32 Heigl, F., Lam, S., Regier, T., Coulthard, I. and Sham, T.K. (2006) ‘Time-resolved X-ray optical luminescence from tris(2-phenylpyridine)iridium’, J. Am. Chem. Soc. Comm.,Vol. 128, No. 12, pp.3906–3907.

33 Rosenberg, R.A., Shenoy, G.K., Heigl, F., Lee, S.T., Kim, P.S.G. and Zhou, X.T. (2005) ‘Determination of the local luminescence sites in ZnS nanowires from X-ray excited optical luminescence’, Appl. Phys. Lett., Vol. 87, pp.253105(1–3).

34 Rosenberg, R.A., Shenoy, G.K., Tien, L-C., Norton, D., Pearton, S., Zhou, X.T. and Sham, T.K. (2006) ‘Anisotropic X-ray absorption effects in the optical luminescence yield of wurtzite nanostructures’, Appl. Phys. Lett., Vol. 89, pp.093118(1–3).

35 Heigl, F., Jürgensen, A., Zhou, X-T., Lam, S., Murphy, M., Ko, J.Y., Sham, T.K., Gordon, R., Brewe, D., Regier, T. and Armelao, L. (2006) ‘Dynamic view on nanostructures: a technique for time resolved optical luminescence using synchrotron light pulses at SRC, APS, and CLS’, SRI 2006, AIP Proc., Vol. 899, pp.1202–1205.

36 Nordgren, J.E., Butorin, S.M., Duda, L.C., Guo, J.H. and, Rubensson, J-E. (2002) ‘Soft X-ray fluorescence spectroscopy for materials science and chemical physics’, in Sham, T.K. (Ed.): Chemical Applications of Synchrotron Radiation, Part l. I, World Scientific, Singapore, pp.517–572.

37 Rubensson, J-E., Mueller, D., Shuker, R., Ederer, D.L., Zhang, C.H., Jia, J. and Callcott, T.A. (1990) ‘Excitation-energy dependence in the L2,3 fluorescence spectrum of Si’, Phys. Rev. Lett., Vol. 64, No. 9, pp.001047.

38 Krause, M.O. (1979) ‘Atomic radiative and radiationless yields for K and L shells’, Phys.Chem. Ref. Data, Vol. 8, No. 2, pp.307–327.

39 Materlik, G., Sparks, C.J. and Fischer, K. (Eds.) (1994) Resonant Anomalous X-Ray Scattering: Theory and Applications, North-Holland, Amsterdam.

40 Hayashi, H., Udagawa, Y., Gillet, J-M., Caliebe, W.A. and Kao, C-C. (2002) ‘Chemical application of inelastic X-ray scattering’, in Sham, T.K. (Ed.): Chemical Applications of Synchrotron Radiations Part II, World Scientific, Singapore, pp.850–908.

41 Tang, C.W. and van Slyke, S.A. (1987) ‘Organic electronluminescence diodes’, Appl. Phys. Lett., Vol. 51, No. 8, pp.605–607.

42 Coulthard, I. and Sham, T.K. (1999) ‘Origin of luminescence from porous silicon: an optical X-ray absorption fine structure (XAFS) study at the Si L3,2- and K-edge’, Solid State Comm.,Vol. 110, No. 4, pp.203–208.

43 Jiang, D.T., Coulthard, I., Sham, T.K., Lorimer, J.W., Frigo, S.P., Feng, X.H. and Rosenberg, R.A. (1993) ‘Observations on the surface and bulk luminescence of porous silicon’, J. Appl. Phys., Vol. 74, No. 10, pp.6335–6340.

1242 T.K. Sham

44 Coulthard, I., Jiang, D-T., Zhu, Y.J. and Sham, T.K. (2000) ‘A systematic study of synchrotron light induced luminescence from porous silicon: implications to morphology and chemical effect dependent electronic behavior’, J. Porous Mater., Vol. 7, Nos. 1–3, pp.165–168.

45 Coulthard, I., Jiang, D.T. and Sham, T.K. (1996) ‘VUV induced luminescence of porous silicon’, J. Electron. Spectrosc. Relat. Phenom., Vol. 79, pp.233–236.

46 Sham, T.K., Coulthard, I. and Naftel, S.J. (2002) ‘Multicore, multichannel detection (MCMD) x-ray absorption fine structure studies of thin films’ in ‘Chemical application of inelastic X-ray scattering’’, in Sham, T.K. (Ed.): Chemical Applications of Synchrotron Radiations Part II, World Scientific, Singapore, pp.1154–1212.

47 Buuren, T., Dinh, L.N., Chase, L.L., Siekhaus, W.J. and Terminello, L.J. (1998) ‘Changes in the electronic properties of Si nanocrystals as a function of particle size’, Phys. Rev. Lett.,Vol. 80, No. 17, pp.3803–3806.

48 Wilcox, J.P. and Samara, G.A. (1999) ‘Tailorable, visible light emission from silicon nanocrystals’, Appl. Phys. Lett., Vol. 74, No. 21, pp.80–82.

49 Zhang, Y.F., Liao, L.S., Chan, W.H., Sammtnakein, R., Lee, S.T. and Sham, T.K. (2000) ‘Electronic structure of silicon nanowires: a photoemission and X-ray absorption study’, Phys. Rev. B, Vol. 61, No. 12, pp.8298–8305.

50 Sham, T.K., Naftel, S.J., Kim, P-S.G., Sammynaiken, R., Tang, Y.H., Coulthard, I., Moewes, A., Freeland, J.W., Hu, Y-F. and Lee, S.T. (2004) ‘The electronic structure and optical properties of silicon nanowires: an X-ray excited optical luminescence and x-ray emission spectroscopy study’, Phys. Rev. B, Vol. 70, pp.045313(1–8).

51 Wolkin, M.V., Jorne, J., Fauchet, P.M., Allan, G. and Delerue, C. (1999) ‘Electronic states and luminescence in porous silicon quantum dots: the role of oxygen’, Phys. Rev. Lett.,Vol. 82, No. 1, pp.197–200.

52 Tohmon, R., Shimagaichi, Y., Mizuno, H. and Ohki, Y. (1989) ‘2.7-eV luminescence in as-manufactured high-purity silica glass’, Phys. Rev. Lett., Vol. 62, No. 12, pp.1388–1391.

53 Stathis, J.H. and Kastner, M.A. (1987) ‘Time-resolved photoluminescence in amorphous silicon dioxide’, Phys. Rev. B, Vol. 35, No. 6, pp.2972–2979.

54 Coulthard, I., Antel Jr., W.J., Freeland, J.W., Sham, T.K., Naftel, S.J. and Zhang, P. (2000) ‘Influence of sample oxidation on the nature of optical luminescence from porous silicon’, Appl. Phys. Lett., Vol. 77, No. 4, pp.498–500.

55 Ö üt, S., Chelikowsky, J.R. and Louie, S.G. (1997) ‘Quantum confinement and optical gaps in Si nanocrystals’, Phys. Rev. Lett., Vol. 79, No. 9, pp.1770–1773.

56 Allan, G., Delerue, C. and Lannoo, M. (1996) ‘Nature of luminescent surface states of semiconductor nanocrystallites’, Phys. Rev. Lett., Vol. 76, No. 16, pp.2961–2964.

57 Batson, P.E. and Heath, J.R. (1993) ‘Electron energy loss spectroscopy of single silicon nanocrystals: the conduction band’, Phys. Rev. Lett., Vol. 71, No. 6, pp.911–914.

58 Eisebitt, S., Lüning, J., Rubensson, J-E., Eberhardt, W., Berger, M. and Arens-Fischer, R. (1996) ‘Quantum confinement effects in the soft X-ray fluorescence spectra of porous silicon nanostructures’, Solid State Comm., Vol. 97, No. 7, pp.549–552.

59 Moewes, A., Kumaev, E.Z., Tse, J.S., Geshi, M., Ferguson, M.J., Trofimova, V.A. and Yarmoshenko, Y.M. (2002) ‘Electronic structure of alkali-metal-doped M8Si46 (M = Na,K) clathrates’, Phys. Rev. B, Vol. 65, pp.153106(1–3).

60 Sun, X.H., Tang, H-Y., Zhang, P., Naftel, S.J., Sammynaiken, R., Sham, T.K., Peng, Y.P., Zhang, Y.F., Wong, N.B. and Lee, S.T. (2001) ‘An X-ray absorption fine structure and electron energy loss spectroscopy study of silicon nanowire at the Si L3,2-edge’, J. Appl. Phys.,Vol. 90, No. 12, pp.6379–6383.

61 Sun, X.H., Wong, N.B., Li, C.P., Lee, S.T. and Sham, T.K. (2004) ‘Chain-like silicon nanowires: morphology, electronic structure and luminescence studies’, J. Appl. Phys.,Vol. 96, No. 6, pp.3447–3415.


62 Tang, Y.H., Sham, T.K., Jurgensen, A., Hu, Y.F., Lee, C.S. and Lee, S.T. (2002) ‘Heavily phosphorous-doped silicon nanowires studied by near edge X-ray absorption fine structure spectroscopy’, Appl. Phys. Lett., Vol. 80, No. 20, pp.3709–3711.

63 Maeda, Y., Tsukamoto, N., Yazawa, Y., Kanemitsu Y. and Masumoto, Y. (1991) ‘Visible photoluminescence of Ge microcrystals embedded in SiO2 glassy matrices’, Appl. Phys. Lett.,Vol. 59, No. 24, pp.3168–3170.

64 Sun, X.H., Wong, N.B., Li, C.P., Lee, S.T., Kim, P.S.G. and Sham, T.K. (2004) ‘Reductive Self-assembling of Pd and Rh nanoparticles on silicon nanowire templates’, Chem. Mater.,Vol. 16, No. 6, pp.1143–1152.

65 Xia, Y., Yang, P., Sun, Y., Wu, Y., Mayers, B., Gates, B., Yin, Y., Kim, F. and Yan, H. (2003) ‘One-dimensional nanostructures: synthesis, characterization, and applications’, Adv. Mater.,Vol. 15, No. 5, pp.353–389.

66 Jenkins, R. and Snyder, R.L. (1996) Introduction to X-ray Powder Diffractiometry,John Wiley & Sons, New York.

67 Mansour, A.N., Cook, J.W. and Sayers, D.E. (1984) ‘Quantitative technique for the determination of the number of unoccupied d-electron states in a platinum catalyst using the L2,3 X-ray absorption edge spectra’, J. Phys. Chem., Vol. 88, No. 11, pp.2330–2334.

68 Bzowski, A., Yiu, Y.M. and Sham, T.K. (1995) ‘Charge redistribution in Au-metalloid intermetallics: a Au L2,3 X-ray absorption study’, Phys. Rev. B, Vol. 51, No. 15, pp.9515–9520.

69 Crespo, P., Litrián, R., Rojas, C.T., Multigner, M., de la Fuente, J.M., Sánchez-López, J.C., García, M.A., Hernando, A., Penadés, S. and Fernández, A. (2004) ‘Permanent magnetism, magnetic anisotropy, and hysteresis of thiol-capped gold nanoparticles’, Phys. Rev. Lett.,Vol. 93, pp.087204 (1–4).

70 Bzowski, A., Sham, T.K., Werneirt, M. and Watson, R.E. (1995) ‘Electronic structure of Au and Ag overlayers on Ru(001): the behaviour of the metal d-band’, Phys. Rev. B,Vol. 51, No. 15, pp.9979–9984.

71 Kuhn, M. and Sham, T.K. (1994) ‘Charge redistribution and electronic behavior in a series of Au-Cu alloys’, Phys. Rev. B, Vol. 49, No. 3, pp.1647–1661.

72 Wertheim, G.K., Dicenzo, S.B. and Youngquist, S.E. (1983) ‘Unit charge on supported gold clusters in photoemission final state’, Phys. Rev. Lett., Vol. 51, No. 25, pp.2310–2212.

73 Schmid, G. (1992) ‘Large clusters and collids. Metal in the embryonic state’, Chem. Rev.,Vol. 92, No. 8, pp.1709–1727.

74 Ohgi, T. and Fujita, D. (2002) ‘Consistent size dependency of core-level binding energy shifts and single-electron tunneling effects in supported gold nanoclusters’, Phys. Rev. B, Vol. 66, No. 11, pp.115410(1–5).

75 Tanaka, A., Takeda., Y., Imamura, M.M. and Sato, S. (2003) ‘Dynamic final-state effect on the Au 4f core-level photoemission of dodecanethiolate-passivated Au nanoparticles on graphite substrates’, Phys. Rev. B, Vol. 68, pp.195415(1–5).

76 Hövel, H., Grimm, B., Pollmann, B. and Reihl, B. (1998) ‘Cluster-substrate interaction on a femtosecond time scale revealed by a high-resolution photoemission study of the Fermi-level onset’, Phys. Rev. Lett., Vol. 81, No. 21, pp.4608–4611.

77 Wilcoxon, J.P., Martin, J.E., Parsapour, F., Wiedenman, B. and Kelley, D.F. (1998) ‘Photoluminescence from nanosize gold clusters’, J. Chem. Phys., Vol. 108, No. 21, pp.9137–9143.

78 Zheng, J., Zhang, C. and Dickson, R.M. (2004) ‘Highly fluorescent, water-soluble, size-tunable gold quantum dots’, Phys. Rev. Lett., Vol. 93, No. 7, pp.77402(1–4).

79 Sham, T.K., Zhang, P. and Kim, P-S.G. (2006) ‘Electronic structure of molecular-capped gold nanoparticles from X-ray spectroscopy studies: implications for coulomb blockade, luminescence and non-Fermi behavior’, Solid State Comm., Vol. 138, Nos. 10–11, pp.552–557.

80 See http://krayos.com/Agen/MagLens.html for technical details.

1244 T.K. Sham

81 Sham, T.K. and Wendin, G. (1980) ‘Screening and configuration-interaction effects in the 5p X-ray-photoelectron spectrum of Th metal’, Phys. Rev. Lett., Vol. 44, No. 12, pp.817–820.

82 Kim, P-S.G.P., Zhang, P. and Sham, T.K. (2004) ‘Reductive deposition of Rh nanostructures on n-type porous silicon: X-ray absorption and X-ray excited optical luminescence’, Langmuir, Vol. 20, No. 11, pp.4690–4695.

83 Zhou, X-T., Sham, T.K. and Lee, S.T. (2005) ‘Fabrication, morphology, structure and photoluminescence of ZnS and CdS nanoribbons’, J. Appl. Phys., Vol. 98, No. 10, pp.024312(1–5).

84 Rosenberg, R.A., Shenoy, G.K., Heigl, F., Lee, S.T., Kim, P-S.G., Zhou, X.T. and Sham, T.K. (2005) ‘Effects of in situ vacuum annealing on the surface and luminescent properties of ZnS nanowires’, Appl. Phys. Lett., Vol. 86, No. 26, pp.263115(1–3).

85 Yamamoto, H. (1998) ‘Zinc sulfide’, in Vij, D.R. and Singh, N. (Eds.): Luminescence and Related Properties of II-VI Semiconductors, Nova Science Publishers, Inc., Commack, p.169.

86 Rosenberg, R.A., Shenoy, G.K., Heigl, F., Kim, P.S., Zhou, X.T. and Sham, T.K. (2005) ‘Direct observation of the local structure of luminescence sites in nanoparticles’, Appl. Phys. Lett., Vol. 87, No. 6, pp.253105(1–3).

87 Zhang, P., Naftel, S.J. and Sham, T.K. (2001) ‘Multichannel detection x-ray absorption near edge structures study on the structural characteristics of dendrimer-stabilized CdS quantum dots’, J. Appl. Phys., Vol. 90, No. 6, pp.2755–2759.

88 Zhang, P., Kim, P.S. and Sham, T.K. (2001) ‘Nanostructured CdS prepared on porous silicon substrate: structure, electronic, and optical properties’, J. Appl. Phys., Vol. 91, No. 9, pp.6038–6043.

89 Zhou, X.T., Sham, T.K., Shan, Y.Y., Duan, X.F., Lee, S.T. and Rosenberg, R.A. (2005) ‘One-dimensional zigzag GaN nanostructures’, J. Appl. Phys., Vol. 97, No. 10, pp.104315(1–4).

90 Nakamura, S. (1998) ‘The role of structural imperfection in GaN based blue light emitting diode and laser diode’, Science, Vol. 281, pp.956–961.

91 Fasol, G. (1996) ‘Room-temperature blue GaN laser diode’, Science, Vol. 272, pp.1751. 92 Gilbert, B., Frazer, B.H., Zhang, H., Huang, F., Banfield, J.F., Haskel, D, Lang, J.C.,

Srajer, G. and de Stasio, G. (2002) ‘X-ray absorption spectroscopy of the cubic and hexagonal polytypes of zinc sulfide’, Phys. Rev. B, Vol. 66, No. 24, pp.245205(1–6).

93 Sun, X.H., Lam, S., Sham, T.K., Heigl, F., Jurgensen, A. and Wong, N.B. (2005) ‘Synthesis and synchrotron light induced luminescence of ZnO nanostructures: nanowires, nano-needles, nanoflowers and tubular whiskers’, J. Phys. Chem. B, Vol. 109, No. 8, pp.3120–3125.

94 Look, D.C. (2001) ‘Recent advances in ZnO materials and devices’, Mater. Sci. Eng. B,Vol. 80, pp.383–387.

95 Huang, M.H., Mao, S., Feick, H., Yan, H.Q., Wu, Y.Y., Kind, H., Weber, E., Russo, R. and Yang, P.D. (2001) ‘Room-temperature ultraviolet nanowire nanolasers’, Science, Vol. 292, pp.1897–1899.

96 Hu, J.Q. and Bando, Y. (2003) ‘Growth and optical properties of single-crystal tubular ZnO whiskers’, Appl. Phys. Lett., Vol. 82, No. 9, pp.1401–1403.

97 Pan, Z.W., Dai, Z.R. and Wang, Z.L. (2001) ‘Nanobelt of semiconducting oxides’, Science,Vol. 291, pp.1947–1949.

98 Yan, H.Q., He R.R., Johnson, J., Law, M., Saykally, R.J. and Yang, P.D. (2003) ‘Dendritic nanowire ultraviolet laser array’, J. Am. Chem. Soc., Vol. 125, No. 16, pp.4728–4729.

99 Kong, X.Y. and Wang, Z.L. (2003) ‘Spontaneous polarization-induced nanohelixes, nanosprings, and nanorings of piezoelectric nanobelts’, Nano Lett., Vol. 3, No. 12, pp.1625–1631.

100 Gao, P.X. and Wang, Z.L. (2004) ‘Nanopropeller arrays of zinc oxide’, Appl. Phys. Lett.,Vol. 84, No. 15, pp.2883–2885.

101 Dai, Y., Zhang, Y. and Wang, Z.L. (2003) ‘The octa-twin tetraleg ZnO nanostructures’, SolidState Commun., Vol. 126, No. 11, pp.629–633.


102 Arnold, M., Avouris, P., Pan, Z.W. and Wang, Z.L. (2003) ‘Field-effect transistors based on single semiconducting oxide nanobelts’, J. Phys. Chem. B, Vol. 107, No. 3, pp.659–663.

103 Kroto, H.W., Heath, J.R., O’Brien, S.C., Curl, R.F. and Smalley, R.E. (1985) ‘C60: buckminsterfullerene’, Nature, Vol. 318, pp.162–163.

104 Iijima, S. (1991) ‘Helical microtubules of graphitic carbon’, Nature, Vol. 354, pp.56–58. 105 Seke, K., Ishii, H. and Ouchi, Y. (2002) ‘Functional organic materials studied using UPS and

NEXAFS’, in Sham, T.K. (Ed.): Chemical Applications of Synchrotron Radiations Part I,World Scientific, Singapore, pp.386–461.

106 Ade, H. and Urquhart, S. (2002) ‘NEXAFS spectroscopy and microscopy of natural and synthetic polymer’ in ‘Chemical application of inelastic x-ray scattering’’, in Sham, T.K. (Ed.): Chemical Applications of Synchrotron Radiations Part I, World Scientific, Singapore, pp.285–355.

107 Tang, Y.H., Sham, T.K., Hu, Y-F., Lee, C.S. and Lee, S.T. (2002) ‘Near edge X-ray absorption fine structure study of helicity and defects in carbon nanotubes’, Chem. Phys. Lett.,Vol. 366, Nos. 5–6, pp.636–641.

108 Tang, Y.H., Zhang, P., Hu, Y-F., Sun, X.H., Fung, M.K., Zhang, Y.F., Lee, C.S., Lee, S.T. and Sham, T.K. (2001) ‘Amorphous carbon nanowires: a precursor to multi-wall carbon nanotubes’, Appl. Phys. Lett., Vol. 79, No. 23, pp.3773–3775.

109 Imamura, M., Shimada, H., Matsubayashi, N., Yumura, M., Uchida, K., Oshima, S., Kuriki, Y., Yoshimura, Y., Sato, T. and Nishijima, A. (1994) ‘C K-edge X-ray absorption near-edge structure of carbon nanotubes’, Jpn. J. Appl. Phys., Vol. 33, pp.L1016–L1018.

110 Kuzuo, R., Terauchi, M., Tanaka, M. and Saito, Y. (1994) ‘Electron energy loss spectra of single-shell carbon nanotubes’, Jpn. J. Appl. Phys., Vol. 33, pp.L1316–L1319.

111 Gruen, D.M. and Buckley-Golder, I. (1998) ‘Diamond films: recent developments’, MRS Bulletin, Vol. 23, pp.16–23.

112 Angus, J.C. and Hayman, C. (1988) ‘Low pressure, mertastable growth of diamond and diamondlike phase’, Science, Vol. 241, pp.913–921.

113 Zhou, X.T., Lai, H.L., Peng, H.Y., Sun, C., Zhang, W.J., Wang, N., Bello, I. Lee, C.S. and Lee, S.T. (2000) ‘Heteroepitaxial nucleation of diamond on Si(100) via double bias-assisted hot filament chemical vapor deposition’, Diamond Relat. Mater., Vol. 9, pp.134–139.

114 Lee, S.T., Peng, H.Y., Zhou, X.T., Wang, N., Lee, C.S., Bello, I. and Lifshitz, Y. (2000) ‘A nucleation site and mechanism leading to epitaxial growth of diamond films’, Science,Vol. 287, pp.104–106.

115 Chakrabarti, K., Chakrabati, R., Chattopadhyay, K.K., Chaudhuri, S. and Pal, A.K. (1998) ‘Nano-diamond films produced from CVD of camphor’, Diamond Relat. Mater., Vol. 7, pp.845–852.

116 Chang, Y.K., Hsieh, H.H., Pong, W.F., Tsai, M-H., Chien, F.Z., Tseng, P.K., Chen, L.C., Wang, T.Y., Chen, K.H., Bhusari, D.M., Yang J.R. and Lin, S.T. (1999) ‘Quantum confinement effect in diamond nanocrystals studied by x-ray-absorption spectroscopy’, Phys.Rev. Lett., Vol. 82, No. 26, pp.5377–5380.

117 Morar, J.F., Himpsel, F.J., Hollinger, G., Hughes, G. and Jordan, J.L. (1985) ‘Observation of a C-1s core exciton in diamond’, Phys. Rev. Lett., Vol. 54, No. 17, pp.1960–1963.

118 Matsumoto, S., Sato, Y., Tsutsumi, M. and Setaka, N. (1982) ‘Growth of diamond particles from methane-hydrogen gas’, J. Mater. Sci., Vol. 17, pp.3106–3112.

119 Zhou, X.T., Sham, T.K., Chan, C.Y., Zhang, W.J., Bello, I., Lee, S.T., Heigl, F., Jürgensen, A. and Hofsäss, H. (2006) ‘Preferable orientation of turbostratic BN basal planes from an X-ray absorption study’, J. Mater. Res., Vol. 21, No. 1, pp.147–152.

120 X-ray calculator, CXRO, Web: http://www-cxro.lbl.gov/.

1246 T.K. Sham

121 Widmayer, P., Boyen, H-G., Ziemann, P., Reinke, P. and Oelhafen, P. (1999) ‘Electron spectroscopy on boron nitride thin films: comparison of near-surface to bulk electronic properties’, Phys. Rev. B, Vol. 59, No. 7, pp.5233–5241.

122 Zhou, X-T., Sham, T.K., Zhang, W., Chan, C.Y. and Bello, I. (2006) ‘Cubic phase content and structure of BN films from an X-ray absorption study’, Anal. Chem., Vol. 78, pp.6314–6319.

123 Zhou, X-T., Heigl, F., Regier, T., Coulthard, I. and Sham, T.K. (2006) ‘Time-resolved X-ray excited optical luminescence from SnO2 nanoribbons: direct evidence for the origin of the blue luminescence and the role of surface states’, Appl. Phys. Lett., Vol. 89, pp.21309.

124 Sham T.K. and Rosenberg, R.A. (2007) ‘X-ray excited optical luminescence: understanding the light emission properties of silicon based nanostructures’, ChemPhysChem, Vol. 8, No. 18, pp.2557–2567.

Nanoparticles and nanowires: synchrotron spectroscopy...

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