Even though two generations of electron microscopists have come to accept that the resolution of their instruments is limited by spherical aberration, three different aberration correctors showing that the aberration can be overcome have recently been built [1-3]. One of these correctors was developed by us specifically for forming small electron probes in a dedicated scanning transmission electron microscope (STEM) [3, 4]. It promises to revolutionize the way STEM instruments are built and the types of problems that they are applied to.

As was the case with the Berlin Wall, when a barrier that was once thought immovable finally crumbles, many of the consequences can be quite unexpected. For STEM, the removal of spherical aberration (Cs) as the main resolution limit is likely to lead to a new paradigm in which:

1) The resolution at a given operating voltage will improve by about 3x relative to today's best. When Cs can be adjusted arbitrarily in a STEM being used for microanalysis or dark field imaging, defocus and Cs are set to values that optimally oppose the effect of the 5th-order spherical aberration C5.

In all scattering experiments some measure is need of the strength of the scattering interaction. The scattering cross section, which has dimensions of area, is a quantity that can be defined for any scattering interactions, irrespective of the nature of the scatterer, or the particle or radiation being scattered. To define a scattering cross section, refer to the geometry of Figure 1. If I0 is the incident number of particles, Is the number of particles scattered through an angle θ with an energy loss ΔE, N is the number of scatterers/ unit volume and t is the thickness of the specimen (or length of the scattering region) then where σ(θ ,ΔE) is the scattering cross section. The product N t represents the number of scatterers per unit area as seen by the incident beam

In electron microscopy all scattering arises from the Coulomb interaction between the incident electron and the electrons or nuclei or the atoms in the specimen.

In the past quarter century low voltage SEM (LVSEM) has been developed for use in a wide range of applications, and it has become the main area of innovation in SEM, and now perhaps more widely in electron microscopy. The intellectual challenge (1,2) has been transformed into a vital tool of modern technology in the semiconductor, polymer and chemical industries, and is widely used in the materials and biological sciences (3, 4). It has been claimed by this author and others that LVSEM is the rational and preferred state of SEM, but unfortunately at the time of the first SEMs only higher voltage systems could be built with the technologies available half a century ago (1). The primary advantage of the SEM remains the use of real bulk specimens with a minimum of potentially invasive preparation. A wide area in the range of square mms to sq inches can contain aperiodic, or otherwise infrequent, events in or on the surface, or in a simple snapped cross-section, and always using relatively stable bulk samples.

Single shell carbon nanotubes are members of the Fullerene family of carbon molecules. Typically, single shell carbon nanotubes measure about 0.7 — 3 nm in diameter and are usually several microns in length. Structurally, they can be thought of as narrow graphite sheets that have been bent around the long axis and joined at opposite edges to form long seamless hollow shells of carbon. Typically, a hemispheroidal cap that contains exactly six 5-rings terminates each end, as shown in Figure 1.

Graphite is known to have an in-plane elastic modulus of ∼1 TPa, one of the highest values known. Consequently, it is expected that single shell nanotubes should be very stiff — a fact that makes them potentially useful in fiber reinforced composites. However, because of their small size, it is impractical to measure their stiffness directly by conventional mechanical means. Recently, we demonstrated that thermal vibrations in freestanding multiwalled nanotubes could be used to estimate their stiffness [1].

Electron Microscopy of surfaces has always been a topic for a few stubborn souls. In addition to the standard headaches of electron microscopes (maintenance, breakdown) one has to handle the additional complications of maintaining a clean enough vacuum to make the results mean anything. (Without vacuum control the structures observed may be due to the contaminants, and hence meaningless.) To complicate things further, in addition to conventional ex-situ sample preparation one has to clean and equilibriate the surface while still retaining thin enough regions for useful study. The final step is to work out which techniques provide surface information and learn how to exploit them. It is appropriate to look at some of the key elements that are needed.

Sample Handling: The key to success with surfaces is not really the microscopy, but instead the sample treatment. The first crucial element is control over the sputter cleaning process.(While simple heating would be less damaging, if one wants to look at something other than silicon sputter cleaning is a must.)

Widespread application of metal oxides in catalysis, gas sensing, and as substrates for thin film growth has stimulated a strong interest in the atomic and electronic surface structure of these materials. The electronic structure of many metal oxide surfaces is characterised, and complicated from the theoretical modelling point of view, by the presence of strong on-site Coulomb repulsion (strong correlations) between valence electrons localised on the metal ions. Additionally, experimental studies have to deal with the difficulties associated with the electrically insulating nature of many of these oxides. We have overcome these problems through the development of novel experimental and theoretical techniques capable of providing structural and electronic information about strongly correlated insulating metal oxide surfaces.

The surfaces of NiO and CoO were investigated through elevated temperature scanning tunnelling microscopy (STM) thereby overcoming problems of low electrical conductivity of the samples.We show atomically resolved elevated temperature STM images of (001) cobalt and nickel monoxide surfaces obtained under similar conditions which show an order of magnitude difference in the atomic corrugation heights.

When I (JMG) began my Ph.D. studies under the supervision of Archie Howie in 1975, he was still licking his wounds over one of the very few apparently incorrect publications[l] of his career. I am proud that our recent work has shown that, far from being wrong, Archie and Lee Rudee were actually 25 years before their time. It was in the early 1970’s that electron microscopy instrumentation reached the performance level that was adequate to resolve the interplanar spacings in elemental materials, such as the semiconductors silicon and germanium. It immediately became of interest to examine the structure of amorphous silicon and germanium, in the expectation that direct imaging could resolve the old controversy between micro-crystallite and random network structural models. Porai-Koschits was the first to suggest that amorphous materials comprised extremely small crystalline domains (microcrystals), in contrast Zachariasen[2] proposed that a continuous random network was a better description of glassy materials.

Disordered or ‘glassy’ materials are of significant technological value. Their structure can be investigated by diffraction (neutron, X-ray, electron) techniques. When the available volume is very small (eg thin films, non-homogeneous structures) only electrons are suitable. Cockayne and McKenzie developed a technique of RDF analysis using energy filtered electron diffraction which gives information on structures of amorphous materials through the reduced intensity function ϕ(s) and the reduced density function G(r). This technique has been applied successfully to many systems.

A major limitation to the use of G(r) for characterisation of amorphous materials arises from the fact that, for electron diffraction, the G(r) measured is a combination of all partial radial distribution functions for each species in the solid. In addition, in practice there are many different models which will equally well fit the data. In order to overcome these difficulties, refinement needs to be done using a realistic starting atomic configuration. For this reason, we have produced starting models using ab-initio molecular dynamics (Car-Parrinello scheme).

Major advances in the field of in situ environmental high resolution electron microscopy (EHREM) probe selective oxidation catalysis directly on the atomic scale. Dynamic gas-solid surface interactions are studied in real-time and under realistic reaction conditions to unravel atomic level insights into active sites and structure-property relationships in vital chemical and technological processes [1-3]. The recent developments include a pioneering approach with the controlled environmental cell (ECELL) facilities permanently mounted inside the EHREM [2]. Accessories have been added for simultaneous structural and compositional analyses of the reactor contents in real-time, and using atomic resolution imaging with transmission electron diffraction and parallel electron energy loss spectroscopy (PEELS). We are now developing innovative experimental methods that include very high temperature studies and combining in the same instrument facilities for both EHREM and environmental-SEM (ESEM), (EHREM-ESEM), with attractive possibilities for studying the science of selective oxidation catalysis.

Alkane Catalysis, Chlorofluorocarbons and Nanotubes:

In the domain of transition metal based oxides, discoveries of fundamental mechanisms underlying selective catalyzation have come from EHREM studies.

The pivotal role played by Archie Howie in the development of many areas of electron microscopy is universary acknowledged.Here I would like to highlight his contribution to the quantitative description of secondary excitations, which was an important influence on the development of Z-contrast imaging in zone-axis crystals. Secondary excitations are those such as x-ray emission which occur following a primary scattering event, in this case excitation of inner shell electrons. The first important concept to be realized by Archie was that dynamical diffraction and channeling are different manifestations of the same physical effect, namely, the multiple scattering of electrons within a crystal. Second was the realization that processes which are localized within the unit cell will show a dependence on diffraction conditions, such as incident beam orientation, and could therefore be described quantitatively using dynamical diffraction theory. Precisely the same theory was used to describe the orientation dependence of cathodoluminescence

The development of the STEM for high resolution imaging was of course due primarily to Crewe and coworkers, with an annular detector to allow efficient detection of elastic scattering over a wide angular range.

Carbon tubes or spheres synthesized by arc-discharge are usually mixed with other byproducts, prohibiting direct measurements of their physical properties by the well established optical techniques because a large quantity of pure specimen is required. Electron energy-loss spectroscopy (EELS) is a unique technique that can be applied to probe the electronic structure of a single carbon tube or sphere. In this paper, the classical dielectric response theory is applied to calculate the EELS spectra acquired from a graphitic carbon sphere at various impact parameters. Graphite is an anisotropic dielectric medium whose dielectric function is described by a tensor. A graphitic carbon sphere is composed of concentric graphitic shells whose dielectric tensor in the spherical geometry, under the local response approximation, is given by (Figure 1)

where are the dielectric function of graphite for an electric field perpendicular and parallel, respectively, to the c axis. In the non-relativistic approximation, the surface excitation is calculated by [1]

EELS in the low loss region of the spectra (< 50eV) provides information on excitations of outer shell electrons and thus the electronic structure of a specimen material which determines its optical properties. In this work dedicated EELS methods for the experimental acquisition and analysis of spectra are described which give improved information about the electronic structure near the bandgap region at a spatial resolution in the range of nanometers. For this purpose we made use of a cold field emission STEM equipped with a dedicated EELS system. This device provides a subnanometer electron probe and offers an energy resolution of 0.35 eV. Application of suitable deconvolution routines for removal of the zero loss peak extracts information on the closest bandgap region while Kramers-Kronig transformation deduces the dielectric properties from the measured energy loss function. These methods have been applied to characterize the optical properties of wide-bandgap materials for the case of group Ill-nitride compounds which are currently the most promising material for applications on optoelectronic devices working in the blue and ultraviolet spectral range.

Crystalline interfaces associated with lattice translations such as stacking-faults (SFs) and twin boundaries may be characterized in the transmission electron microscope (TEM) by oscillatory fringe contrast known as SF, or rigid-body-translation, fringes. Such fringe contrast is usually obtained using weak-beam bright-field (BF) or dark-field (DF) conditions with ideally only one reflection, g, excited. The nature of the SF contrast has been explained by dynamical theory of electron diffraction in terms of the Howie-Whelan equations [1]. The fringe contrast arises from an interference caused by a phase shift the diffracted electron beams, undergo upon propagating across the interface. The phase shift, α, is proportional to the projection of the translation vector, R, onto the diffraction vector, g (i.e. α # g.R). The direction of the translation state can be determined explicitly from the SF contrast in a manner analogous to the analysis of dislocation Burgers vector.

An analysis of SF fringes with respect to diffraction theory permits the magnitude of the lattice translation to be evaluated in experimental images.

Since heterogeneous catalysis depends on surface chemistry, insights into the structure of supported catalytic materials is vital for understanding catalysis. We have developed a mass spectroscopic technique using very high angle annular dark field images in a scanning transmission electron microscope (STEM), that gives the number of atoms per cluster. We have also developed a robust interactive computer program to analyze these images rapidly. In this proceeding, we will present our method and results on analyzing ultra-small and dispersed Pt clusters on amorphous carbon, as well as our preliminary results of Pt on γ-A12 O3.

The ultra-small Pt clusters on the amorphous carbon Cu grid were made by a proprietary technique. Imaging for the STEM-based mass-spectroscopic technique was performed on a Field Emission Gun (FEG) VG HB601 STEM operated at l00kV. The absolute measured intensity from the clusters were converted to scattering cross-sections, which can then be converted to number of atoms: More details of this method can be found in Singhal, Yang and Gibson.

Introduction

The morphology of Si surface has much effect on physical and chemical properties of the surface. It is well known that the structure of Si surface can be modified by deposition of metals, and extensive studies have been carried out on the modified surface structures of Si. These studies were carried out on clean surfaces of Si in an ultra-high vacuum (UHV). However recently, it has been shown by Kamino et al.[l] that molten Au atoms remove a surface amorphous layer on the Si surface, making the surface clean even in a non ultra-high vacuum. The present study is an extension to the previous one (Kamino et al.) and Si surfaces modified by-molten Al atoms in a conventional vacuum were observed by means of TEM.

Experimental

The specimen used in this study was a mixture of pure Al particles and Si particles, the latter being prepared by crushing a Si single crystal into pieces.

INTRODUCTION

It has been reported by many authors that metastable β-Ti alloys such as Ti-Nb, Ti-V and Ti- Mo indicate the negative temperature dependence of electrical resistivity at low temperature since Ames et al., found the dependence on Ti-Nb alloys in 1954. Williams et al., suggested that the phenomenon resulted from formation of athermal co-phase crystals. In order to clarify the origin, electron diffraction pattern analyses were carried out using a specimen cooling stage and both of start and finish temperatures of formation of athermal co-phase crystals were decided. However, it is necessary to reveal the morphology of athermal ω-phase crystals to study the relationship between the electrical resistivity change and appearance of athermal ω-phase crystals. On the other hand, although electron diffraction pattern analyses are able to identify the atomic structures of crystals, which appear the diffraction spots, it is not so easy to reveal the detailed morphology of the crystals. Also, it has to consider effects of electron irradiation to transformations.

Bimetallic alloy particles supported on high-surface area carbon are widely used as electrode materials. Because of their stability, Pt-Sn alloys are of particulate interest for a variety of applications. Although the reasons for the superior properties of these catalysts are not fully understood, the dispersion, structure, and compositional homogeneity of the alloy clusters, are clearly important factors in determining their catalytic activity. In order to identify the composition and structure of Pt-Sn base nanoparticles, various characterization techniques, mainly X-ray diffraction (XRD) and transmission electron microscopy (TEM) [1], have been applied. However, very often conflicting conclusions regarding to the structure of Pt-Sn particles have been reported. The focus of the present work is to examine a commercially available carbon supported bimetallic Pt-Sn catalyst, with a nominal Pt:Sn wt. ratio of 2:1 (500N and 900N samples) and 3:1 (E270 sample), using high resolution electron microscopy (HREM), electron and X-ray diffraction techniques, and microchemical analysis by energy dispersive X-ray spectroscopy (EDS).

We are developing a new family of heterogeneous catalysts for hydrogenation catalysis. Catalyst synthesis is accomplished using colloidal polymerization chemistry which produce high surface area xerogel catalysts. These xerogels have been synthesized by one-step sol gel chemistry. These catalysts contain ruthenium and modifiers such as gold occluded or incorporated in a titanium oxide matrix. The materials, especially the modified systems exhibit favorable performance in microreactor evaluations for hydrogenation reactions and exhibit high activities. Nanostructural studies have revealed that the materials contain dispersed catalyst clusters which are desirable microstructures for the catalysis since the majority of the atoms are exposed to catalysis and are potentially active sites.

The composition and atomic structure of the xerogel catalysts containing ruthenium and other metals have been examined using our in-house developments of environmental high resolution electron microscopy (EHREM) the atomic scale [1-3] and low voltage high resolution SEM (LVSEM)[4] methods.

A solid-liquid interface in the Al–Si system has been observed at near-atomic resolution by in-situ heating experiments inside transmission electron microscopes. Chemical mapping was also attempt to detect distribution of constituent atoms near a solid–hquid interface.

Mixtures of Al particles and Si particles, the diameter of which ranged from 200nm to 800nm, were mounted on a specimen–heating holder developed by Kamino and Saka and examined in a Hitachi H–9000NAR and a Hitachi HF–2000 electron microscope, operated at accelerating voltages of 300 and 200kV, respectively. Heating these specimens above the melting point of pure Al inside the microscopes resulted in the melting of the Al particles. The liquid Al reacted with nearby Si, leading to the formation of Al–Si alloy phase. An interface between solid Si and Al–Si alloy hquid and also an interface between solid Al and Al–Si alloy liquid were observed. The dynamical behaviour was recorded continuously using TV–VTR systems. The distribution of Al and Si was visualized using a Gatan imaging filter GIF analyzer.

Small atomic clusters are of great importance for applications such as catalysts whose activity depends on the surface of the cluster. Attempts to determine the atomic short-range order and size of clusters have been made by analyzing the extended X-ray absorption fine structure (EXAFS). However, the analysis was made on an average of many small clusters. Analysis of extended energy-loss fine structure (EXELFS) in an electron energy-loss spectrum (EELS) has developed to the point where in some cases, the quality of the results is comparable to its X-ray analogue, EXAFS. No other technique provides nanometer-scale spatial resolution of the analyzed area for determining the atomic structure. Most EXELFS analysis has been performed on the K-ionization edge of lighter elements. For heavier elements, a more complex ionization edge such as the L-edge has to be used, due to the inefficiency of collecting high quality EEL spectra at higher energy-losses (Z > 18).