Approaching atomic-resolution electron microscopy

Ultramicroscopy 18 (1985) 63-76 63 North-Holland, Amsterdam

APPROACHING ATOMIC-RESOLUTION ELECTRON MICROSCOPY

David J. SMITH *, R.A. CAMPS, L.A. FREEMAN, M.A. O'KEEFE ** , W.O. SAXTON and G.J. WOOD * High Resolution Electron Microscope, University of Cambridge, Free School Lane, Cambridge CB2 3RQ, UK

Received 1 July 1985; presented at Symposium January 1.985

The Cambridge University 600 kV high-resolution electron microscope, supported by local facilities for image simulation and processing, has been successfully used to characterise the microstructure of a wide range of materials atresolutions on, or approaching, the atomic scale. After short descriptions of the microscope and computer facilities, attention is first directed towards the problems of imaging artefacts and beam misalignment. The usefulness of Weak-Phase-Object images for preliminary assessments of image detail, such as dumbbells and possible oxygen atom resolution, is stressed. It is shown that the well-known imaging artefact, namely dumbbells, is an interference phenomenon which can be generated even when the crystal being imaged does not have the classic "atom-pair" projected structure. Methods for assessing resolving power and the validity of information beyond the resolution limit are considered. Particular applications to the study of surface rearrangements, nonstoichiometric ruffle and defect annealing in cadmium telluride are briefly described.

1. Introduction

The achievement of atomic resolution in the electron microscope not only implies proper electron-optical design and adequate instrumentation but it also requires a thorough understanding of the imaging process. During optimisation of the Cambridge University 600 kV High-Resolution Electron Microscope (HREM), improvements in instrumentation and concurrent developments in image simulation and processing facilities were necessary before atomic-resolution capabilities could he approaclied, and utilised, on a routine basis. In this paper we summarise some of these advances as well :as review the wide range of materials which have been studied with this particular microscope. These examples serve to demonstrate the increasing impact that atomic-resolution studies are having on our knowledge and

* Present address: Cefiter for Solid State Science, Arizona State University, Tempe, Arizona 85287, USA.

~* Present address: National Center for Electron Microscopy, Lawrence Berkeley Laboratory, University of California, Berkeley, California 95920, USA.

understanding of materials' properties. The design philosophy of high-voltage HREMs

is straightforward. In the absence of workable aberration-correction devices, the most plausible way of overcoming the aberration-wavelength limitations on resolution is to increase the accel- erating voltage [1]. This potential has been appre- ciated for many years [2], but it has only been comparatively recently that technical improvements have allowed these possibilities to be realised in a number of laboratories, though rarely on a routine basis (see ref. [3] for details of these pro- jects). The goal of atomic resolution was the motivation for a major research and development project at Cambridge University and details of the design, construction and commissioning of the 600 kV instrument have been given elsewhere [4-8]. Our subsequent objective, which is the main topic of this paper, has been to establish this microscope as a high-resolution facility and then to apply it to a number of promising materials where it was anticipated that its extreme resolving power of better than 2 A would permit elucidation of structural defects on the atomic scale. The indispensable role of the image simulation/processing facilities

0304-,3991/85/$03.30 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

64 D.J. Smith et al. / Approaching atomic-resolution electron microscopy

to many of these high-resolution studies will also be emphasised.

2. Instrumentation

2.1. The basic microscope

The 600 kV microscope - shown schematically in fig. 1 - basically consists of the AEI EM7 microscope column and a high-voltage power supply and air-insulated four-stage electron accelerator supplied by Haefely, although there have been many local additions and refinements which have been instrumental in reaching the design performance. The column bolting was strengthened so that it could be suspended from above and the standard lens supplies were replaced by higher-stability versions. Modifications to the feedback cir- cuits and cable layout improved the original stability, drift and ripple of the high-voltage generator (typically 1 ppm levels are required for atomic resolution).

The most crucial of the on-site developments, and certainly the most novel, is the three-point

Fig. 1. Outline sketch of the Cambridge University 600 kV high resolution electron microscope: G = L a ~ electron gun; A electron accelerator; M zmicroscope eolunm; S=support girders; P = pneumatic cylinder.

suspension system with which the entire instrument is supported on the walls of the microscope room. This system incorporates pneumatic cylin- ders and rolling rubber diaphragms with self- levelling facilities, giving resonant frequencies of about 0.5 Hz. Ambient vibrations, monitored by a piezo-electric accelerometer device, are attenuated by factors in excess of 100 to levels where the effects of nearby machinery, or passing traffic, on operating performance are minimal at any time of day.

The electron gun of the microscope uses an indirectly heated lanthanum hexaboride rod cathode [9]. The lower work function of LaB 6 leads to an order of magnitude increase in bright- ness relative to the conventional tungsten hairpin, so that illumination with smaller angles of inci- dence (i.e. better spatial coherence), which is vital for highest resolution, can be obtained. At the same time, higher operating magnifications are also possible allowing visual correction of image astigmatism directly from the fluorescent screen if required. An added benefit of using LaB 6 is that rod lifetimes are generally well in excess of 300 hours.

Construction of a specimen stage and holder adequate for the demands of atomic-resolution microscopy was carried out by Mr. J.H. Lucas (Rickling, Essex, UK). The stage itself rests on the lower objective lens polepiece for mechanical and thermal stability, and a side-entry detachable load- ing mechanism is provided for insertion of the holder into the microscope. The goniometer (tilting) holder provides two orthogonal tilts of + 30 ° without loss of positional stability relative to non- tilting holders, as demonstrated in fig. 2 which shows an image of lattice fringes having a separation of only 0.72 A. The visibility of this lattice detail is also a sensitive indicator of the overall mechanical stability of the microscope.

The objective lens for the microscope column has been modified substantially to improve its imaging performance. With an 11 mm polepiece gap and 7 mm bores, and with the high permeabil- ity alloy Permendur to minimise saturation, calculations indicate values for the spherical, Cs, and chromatic, C c, aberration coefficients of Cs = 2.48 mm, Cc = 2.82 mm, with the specimen located at

D.J. Smith et al. / Approaching atomic-resolution electron microscopy - 65

Fig. 2. Lattice fringes of 0.72 A spacing from a small particle of silver imaged at 575 kV.

the centre of the polepiece gap [10], which would give intuitive image resolutions extending out to better than 1.8 A [8].

An image pickup and display system has been attached to the base of the microscope [11]. It uses an Image-Isocon tube (type P8041), with a 40 mm photocathode, in direct contact with a phosphor- coated fibre-optic glass-plate which also acts as a vacuum seal. The output permits observation at current densities well below that required for normal fluorescent screen viewing; it also allows image processing and control of critical microscope parameters by a dedicated on-line minicomputer [12]. The observation of beam-sensitive specimens is facilitated and dynamic events occurring at the atomic level can be observed and documented by recording on video-tape.

2.2. Modifications and improvements

As with almost any "prototype" instrument, unexpected factors proved just as critical in limit- ing the performance of the 600 kV microscope as did any initial failure to met the required toler- ances for the mechanical and electrical stabilities. These various shortcomings had first to be pin- pointed, and then overcome, before atomic resolution could be achieved. Vibrational turbulence due to the lens and pump cooling water, and defective resistors in the high-voltage bleeder chain were

reasonably straightforward to locate. More subtle were the stray AC magnetic fields originating from faulty mains supply wiring in nearby buildings, and a careful, sometimes tedious, process of elimination was needed to isolate each source. As a further precautionary step, it was also established that the final image was most sensitive to the stray fields in the vicinity of the first inter- mediate image, so mumetal shielding was installed at this level inside the coltmm at the small sacrifice of selected-area diffraction capability. Finally, the installation of facilities for monitoring the high- voltage stability, mechanical vibrations and the external AC fields has proven well worthwhile. Experience has shown that these factors are the most likely to develop faults, and the monitors enable these to be rapidly located and diagnosed.

Although the resolving power of an electron microscope should be limited by fixed instrumental parameters, in practice the resolution actually achieved depends as much on the skill of the operator in setting up proper imaging conditions within the objective lens. Not only must accurate focus settings be chosen but the incident electron beam must be closely aligned with the optic axis and the image must be as free as possible of residual astigmatism. The stringent requirements on these parameters tax even the most experienced operators so that the literature of high-resolution electron microscopy is full of examples where in- correct imaging procedures have been used [13]. Fortunately, it was possible to connect the output signal of the image pickup system into a suitably designed framestore (a GEMs MkI) allowing image assessment to be carried out automatically using a CA 4/90 minicomputer, followed by appropriate adjustments to the focusing, stigmating and alignment controls [14]. Criteria based on minimising contrast were used and it has been found that the final correction on-line can be about twice as accurate as an operator can routinely achieve [12,14]. Measurement of the optical diffractogram shown in fig. 3, for example, indicated that a residual astigmatism of less than 50 A had been achieved after two iterations of the on-line correction system.

A further device has been attached to the microscope which enables accurate beam alignment


material can sometimes relax this requirement i Reference should be made to recent reviews elsewhere for further details [17,18]: we restrict our discussions here to the particular problems of imaging artefacts and the effects of beam misalign- merit, whilst we consider resolution criteria and the validity of information beyond the resolution limit in a later section.

3.1. Image simulation and processing

Fig. 3. Optical diffractogram from a 500 kV micrograph of amorphous carbon recorded after on-line computer correction of objective lens astigmatism and incident beam misalignment.

to be achieved in the absence of on-line computer facilities [8,13]. It is a variant of the method described by Zemlin [15], and relies on recognition of differences in image appearance when the incident illumination is tilted by equal and opposite amounts in either of the two orthogonal tilt direc- tions. The mean position is continually adjusted until the image texture and appearance is identical at both tilted positions. With practice, this method can be carried out as quickly and reliably as astigmatism correction.

3. Image interpretation

The appearance of a high-resolution electron micrograph depends, in a complicated and non-intuitive manner, on characteristics of the object and on the properties of the imaging lenses. For- tunately, electron diffraction theory is well-ad- vanced, at least for thin objects, and computer programs based on the multi-slice formulation [16] aCe generally available which can replicate the electron scattering, as well as the imaging process, which occurs inside the HREM. Image simulations should normally be regarded as an essential part of interpreting images of structural imperfections, especially for details near the limits of resolution, although previous experience with a certain

Whilst it is possible to assess imaging conditions with an optical bench, and rely on access to a central computer for image-simulation capabilities, our experience, also reported by others [19,20], has been that a dedicated local minicomputer can help immeasurably in many ways. For example, structural refinements can be expedited by interactive model-building using comparisons of simulated images with digitised micrographs. Similar comparisons, for through-focal series of.images, can be used to quantify microscope imaging parameters [21] or, alternatively, the contrast transfer characteristics can be established by digital analysis of diffractograms computed from images of amorphous objects.

The present generation of "super" minicom- puters, such as the: SEL 32/27 installed in our laboratory, are able to handle the computational demands in respectable periods, although Fourier transformations of large arrays sizes can still take considerable time (about 2½ min for a 512 x 512 image in our case). Interfaces were arranged lin- king a small CA 4/10 minicomputer to the SEL host, to a (second) GEMS framestore, and to an existing microdensitometer and filmwriter; a large resident program was derived for the CA to allow it to manage the peripherals in response to instruc- tions from the host or from a small keypad/track- erball beside the user terminals [22]. The combined system allowed a variety of interactive image processing tasks to be accomplished with the previously established Semper system [23], in parallel with multi-slice image simulations using the SHRLI programs [18] and others, on a time-sharing basis.

Amongst other things, careful evaluation of simulated images has been used to identify i the nature of stacking defects in a martensitic CuZnA1

D.J. Smith et al. / Approaching atomic-resolution electron microscopy 67

alloy [24], and of a glide opdation responsible for a twin boundary in SnO 2 [25]; to confirm the direct interpretability of HREM images of a AuMn alloy [26] and of polytype stacking sequences in SiC [27]; to demonstrate that the images of aromatic hydrocarbons are relatively insensitive to thickness and defocus (because of the relatively weak high-order scattered beams) [28]; and to explore the Fresnel contrast developed in bright- and dark-field conditions around a twin boundary in Cu [29]. A substantial modification to an existing multi-slice program allowed convergent-beam diffraction (CBED) patterns to be simulated for crystal thicknesses up to several tens of nanometres [30].

3.2. Imaging artifacts

It is well established that the transfer characteristics of an objective lens vary rapidly with the defocus thereby causing the final image appearance to alter drastically, especially for large- unit-cell materials when there are many diffracted beams scattered throughout reciprocal space. Beam misalignment, or residual objective lens astigmatism, can also contribute to artefactual image detail. Moreover, contrast reversals occur periodi- cally for small-trait-cell materials, producing Four- ier, or "self", images [31] and it can then be very difficult, if not impossible, to recognise the correct focus'for direct image interpretation without reference to a nearby defect or crystal edge. A further complication for small-unit-cell materials is the occurence in some projections of the so-called "dumbbell" or "atom-pair" images which usually occur in the vicinity of the first extinction contour where the transmitted beam is of low intensity and the image is dominated by second-order effects. We consider it worth re-emphasising, because of continuing confusion in the literature, that these dumbbells are interference phenomena which re- flect both the crystal structure and the instrumental resolution but they are not necessarily related to any real atomic column positions.

For all these reasons, it is highly recommended that any quantitative ~gh-resolution image interpretation should rely on a foundation of multi-slice computation of a through-focal, through-thickness

series of images, as well as a through-resolution series of weak-phase-object (WPO) images. In the latter, images are simulated with progressively more diffracted beams contributing to the image but all with unit transfer (i.e. without including any microscope or CTF damping effects). Such WPO images are independent of lens parameters, are very easy to compute and digest, and give a very good representation of the thin crystal image near optimum defocus [27,36].

A typical example for the ferroelectric material BaTiO3 imaged in the (001) projection is shown in fig. 4. The top-left image depicts the projected atomic positions within the cell used for the simulations. At a WPO resolution of 3.99 A, when only 5 beams contribute to the image, the cation positions are already resolved. With better resolution, the image appearance changes rapidly, although retaining its symmetry, and at a WPO resolution of 1.79 A it seems that the oxygen atoms are also being resolved. However, this is clearly an artefact

Fig. 4. A Weak-Phase-Object (WPO) series of image of BaTiO3 showing the effect of increasing resolution on the visibility of sub-unit-cell detail. Note that a WPO resolution of 0.89 ,k is required before the individual oxygen atoms are resolved.

68 D.J. Smith et aL / Approaching atomic- resolution electron microscopy

since the features of the image alter again as further beams are used, and is it not until a WPO resolution of 0.89 A that the image "settles down" and could be considered to be a true representation of the object; i.e., these novel results suggest that oxygen atoms in BaTiO3 will not be resolved in any existing HREMs.

It is interesting to compare these WPO images with full multi-slice simulations. Fig. 5, for example, shows calculations for a 1000 kV HREM having an objective lens spherical aberration coef- ficient of 2.6 ram. The top row is for a defocus of -450/~ , which is close to the optimum, and t h e bottom is for - 900 A, with increasing thicknesses from left to right. In the thinnest crystal image, at the smaller defocus, there are black spots at the cation column positions and it would be reasonable to suppose, in the light of the WPO images, that these are being imaged although the oxygen columns are not. It is remarkable, but probably no more than coincidental, that the image at the larger defocus and a crystal thickness of 168 A has almost exactly reverse contrast when compared with the WPO image having a resolution of 1.79 A.

The WPO images also provide an interesting basis for discussing the "dumbbell" effects, and fig. 6 shows various members from the WPO series of (a) tetragonal tin dioxide, SnO2, imaged in (100), and (b) diamond cubic silicon imaged in (110). Note that what these two materials have in common is a hexagon-shaped electron diffraction pattern, although only Si has the classic projected

Fig. 5. Simulated images of BaTiO 3 for a 1000 kV, C s = 2.6 ram, HREM showing effect of increasing specimen thickness (as marked). Top row for defocus of -450 A; bottom row for -900 A.

Fig. 6. WPO (increasing resolution) series of images for (a) SnO2 in (100) and (b) Si in (110). Note that a WPO resolution of 1.36 A is needed to resolve the Si atom-pairs whilst 1.41 A is needed to resolve the O atoms in SnO2.

atom pair. The silicon atom pairs are "resolved" at a WPO resolution of 1.36 A, and the oxygen atoms in SnO 2 at 1.41 A. Contrast these simulations with the 500 kV experimental micrographs shown in fig. 7, which are from the vicinity of the first thickness extinction contour. For silicon, the white spots ( - 1 . 8 A apart) appear almost at the positions of the atom pairs. For tin dioxide, the white spots (separation of - 1.5 A) are again close to the positions of the oxygen atom columns ( - 1.8 A apart). Fig. 7c shows a "dumbbell"-like image in Au, but note that there are no atomic columns to "match" the white spots! Our overall conclu- sion is that it would be preferable, and far less confusing, to label those effects as interference lattice images, rather than atom pairs, and thereby avoid any implication that the spots are related in any simple manner to crystal structure [3,27].

3.3. Effects of beam misalignment

It has only been comparatively recently that high-resolution electron microscopists have really begun to appreciate the consequences of incident beam misalignment [13,32]. Previously, it had generally been assumed that the current or high-voltage centre methods would be adequate. However, with the increasing trend to operate with higher-excitation objective lenses, rotational symmetry of the lens field is no longer retained and the true (coma- and dispersion-free) axis is dis- placed. Images recorded under such conditions are again likely to contain artefacts which are unre-


Fig. 8. Simulated optical diffractograms at the optimum (Schemer) defocus with increasing beam tilt (as marked). Even with 2 mrad tilt, the visual effect on the ODM is marginal.

Fig. 7. Experimental 500 kV micrographs of (a) Si in (110); (b) SnO 2 in (100); (c) Au. All show second-order "dumbbell" effects due to being recorded at the first thickness extinction contour.

lated to the features of the specimen. Unfortunately, as shown in fig. 8, the magni-

tude and direction of misalignments of 1-2 mrad cannot be easily detected using the standard optical diffractogram from the image of an amorphous object since its basic effect when small is only to alter the phase of the transfer function of the

objective lens. If k 0 measures where the mis- aligned beam intersects the Fourier plane, the additional factor in the transfer function has the form

e x p [ - 2i(CsXZk 2 - D ) h 2 k • k0)] (1)

for a specimen periodicity k and underfocus D. In reduced units, the actual phase shift q~ caused by off-axis misalignment can be written [13]

~ = - 2 - ( k 2 - O ) k . k o . (2)

The antisynunetric phase shifts cause lateral displacement of fine image features, and destroy any centrosymmetry otherwise to be expected in the image although translational symmetry is unaffected by misalignment. In most circumstances, crystal misalignment by similar amounts has a small effect only, though anomalous effects can be observed for images of thick crystals, particularly near extinction contours in projections containing dynamically forbidden reflections [33].

Because of the image sensitivity to beam misalignment, point-group assignments should never be made on the basis of high-resolution images alone, and a supplementary technique such as convergent-beam electron diffraction should then be employed. Beam misalignment should be eliminated, wherever possible, by one of the methods described above, preferably just before recording a micrograph, as should already be the custom for adjusting image astigmatism.


4. Assessment of performance

While it would be simplistic to suppose that there is any particular resolution distance below which image detail is undetectable and above which it is entirely reliable, and while what is resolved by a given image is to some extent necessarily a function of the a priori information available about the specimen, it is nevertheless useful to attempt to quantify the resolving power of an instrument - if only for the purpose of monitoring or comparing its level of performance. We examine here two different approaches to practical assessment of the resolving power, and the feasibility of measuring distances or locating atomic positions with an accuracy far beyond the resolving power.

4.1. Resolution criteria

Two criteria are widely used for the purpose of defining resolution, namely the spatial extent of diffractograms from images of thin amorphous objects, and the detailed comparison of experimental micrographs with an atlas of images simulated for a range of instrumental parameters.

The diffractogram, traditionally an optical diffractogram, but increasingly a digitally calculated diffractogram, reveals the squared modulus (intensity) of the instrumental transfer function, since the specimen transform itself is presumed to be diffuse, isotropic and relatively slowly varying. Its outer limit identifies at once the finest object period transferred linearly to the image. The illumination parameters (divergence and effective focal spread) can be identified by careful evaluation of the diffractograms of a focal series [34] as can the more obvious parameters of focus and astigmatism, though not, as noted above, small levels of beam tilt. It is in practice, however, a highly subjective decision at what spatial frequency the diffractogram signal fades entirely into the diffuse background arising from recording noise in the image. The practice of recording a pair of immediately consecutive images allows this decision to be made more objective, either by super- posing the images slightly out of register before calculating the diffractogram [35], or by comparing the Fourier transform in a radial spatial frequency correlation function [36]. It is almost

never safe to rely on the highest-order reflection visible in the diffractogram of an image of an inorganic crystal for the purpose of assessing the useful resolution of an image, as second-order contributions to the image intensity, some of them wholly unaffected by focus spread or illumination divergence, often extend well beyond the limit of linear transfer [52].

Given a well characterised specimen such as the block oxide Ti2Nb10029, the illumination parameters may also be derived with reasonable accuracy by comparing recorded .images with a set of through-focal, through-thickness image simulations for a variety of different illumination parameters [21]. For large unit cell materials at least (i.e. those with densely populated Fourier space), a simple assessment of the useful resolution in the image may be obtained by comparing optimum focus images with a through-resolution WPO series (see above) [37]. Image matching has rarely been quantitative in the past, and our own practice has not been exceptional in this respect; it is likely that greater sensitivity will be possible in future by regular combination of image processing and simulation to allow quantification of the match.

4.2. Beyond the resolution limit

The reliable interpretation of image detail at or beyond the resolution limits of a particular microscope is an activity which should always be accom- panied by image simulations. However, there are some circumstances for which much finer specimen information can be obtained. One of these relates to rigid b o d y displacements at edge-on boundaries. Provided that there is adequate signal- to-noise to define lattice fringe positions accu- rately on either side of the border, then it is possible to measure relative displacements normal to the boundary with an accuracy approaching 2% of the lattice parameter [29,38], even though there may be little information available about the actual atomic column locations in the same micrograph. In this case, the accuracy achieved depends on the regularity of the extended sample. Such an assumption is likely to be invalid in the immediate vicinity of structural defects such as dislocations and grain boundaries. A detailed analysis of


specific model situation has been described [39]. It appears that the relative positions of atomic col- urnns can probably be located in experimental micrographs to within one-sixth of the resolution limit of the particular HREM and, by using object restoration by image processing, that better than about 5~ is sometimes possible. Image averaging, on the other hand, without compensation for the phase-shifting effects of the contrast transfer function, can sometimes approach this sort of accuracy [40] but the method is inapplicable to aperiodic objects.

5. Applications

The great attraction of high-resolution electron microscopy over other techniques for characteris- ing materials arises from its ability to provide localised real-space information on the atomic scale about structural defects and thereby give unique insight into many materials' properties. Optimis- ation of the 600 kV instrument, including the routine incorporation of an accurate alignment procedure, has enabled this potential to be realised on an effectively daily basis and a wide range of different materials from local, national and over- seas sources have subsequently been studied to good effect with this machine. Some of the more interesting and significant results are briefly described here: reference should be made to papers published elsewhere for further details of applications and information about other materials [4-8,24-29,41-51].

ble. By allowing etching conditions to develop inside the microscope, surface carbonaceous materials could be removed, and fig. 9 shows the edge of a small gold particle which had partially reconstructed into the 2 x I (110) "missing-row" structure [42]. Subsequently, further observations were undertaken using extended gold surfaces rather than particles, in what has been termed a surface profile imaging mode, and these have provided novel and wholly unexpected results (see ref. [43] for full details and other references). The (100) surface was often found to contain dislocations, but only in the presence of surface steps, which apparently acted as asymmetric potential barriers to atomic motion, was there any substantial surface bulk rearrangement. The extended (111) surfaces developed a buckling which led eventually to a pronounced hill-and-valley morphology. Finally, the extended (110) surfaces were generally found to be highly mobile and microscopically rough. The movement of entire atom columns between successive exposures, as shown in fig. 10, was a common phenomenon observed to take place on this surface.

5.2. Nonstoichiometric rutile

Novel structural features have been established by recent HREM observations of slightly reduced rutile, TiO2_ x (0 < x < 0.0035). These include the absence of crystallographic shear planes (CSP) in quenched materials, the existence of {100} platelet

5.1. Atomic imaging of small particles and surface rearrangements

The most striking observations made with the 600 kV instrument, with important long-term im- plications for catalysis and surface science research, originated through the study of multiply- twinned particles of silver and gold. Detailed observations, summarised in [41], confirmed, for the first time, the presence of dislocations and stacking faults in particles less than 100 A in diameter, and it was also realised that atomic level information about facetting and surface steps was availa-

Fig. 9. Region of a small Au particle, imaged in a (110) projection, showing a partial 2 × 1 "missing-row" reconstruc- tion of the (110) surface [42].


Fig. 10. Successive 500 kV images (time separation 15 s) of an extended (110) Au surface showing the movement of entire atomic columns (arrowed).

defects, and lateral and longitudinal disorder in the CSP fine structure [44]. On the basis of these results, new structural models were developed for both cation interstitial and anion vacancy defects [45], and these models should account for various point defect phenomena associated with rutile, and are likely to be applicable to other nonstoichiometric oxide systems. Samples of chromia-doped rutile, Ti(Cr)O1.92, indicated the presence of exten- sive disorder along the CSP despite the contrary evidence of the electron diffraction pattern which predicted a well ordered superstructure [46]. Image simulations of this material also established that, provided the microscope had an interpretable resolution limit of better than 2 A, then the cation arrangements along the CSP could be read off directly with, in particular, the so-called face- shared (C) and anti-phase (A) sites being dis- tinguishable [46].

Attempts have been made to study the individual small defects responsible for the nonstoichiometric phase TiO2_ x but only small groups o r clusters of these defects have so far been imaged [47]. Indirect evidence for the predominant nature of the small defects has been obtained from the magnitude and sense of the lattice displacement at a CSP termination [48]. The extrinsic displace-

Fig. 11. Termination of an extended crystallographic shear defect in TiO1.9666 [48]. The closure failure R = [0,0.5,0-.5] established the extrinsic nature of the defect.

ments seen in rutile, such as shown in fig. 11, support the cation interstitial model [48], whereas it is interesting and significant that intrinsic vec- tors have been observed in WO 3_x [49].

5.3. Defect annealing in cadmium telluride

Real-time vieWing and recording of dynamic processes, including defect rearrangements, inside the HREM at the lattice or atomic-resolution level should, in principle, provide enhanced insight into the reaction mechanisms of many solid state transformations. This" potential was demonstrated in our observations of the semiconductor compound cadmium tdluride [50]. Under the influence of the incident electron beam, defect annealing was in-


: o

duced to taken place and several types of phenomena were documented using the videotape recording facilities. Since the projected CdTe pairs of atomic columns, with spacing of 1.68 .~, could not be separately resolved, these events were recorded at the lattice resolution level. Nevertheless, some of the microstructural changes observed were directly interpretable in terms of atomic motion. These included: (i) gradual disappearance of a (111} planar defect, as shown in fig. 12; (ii) changes in length of extrinsic and intrinsic stacking faults; (iii) elimination of an intrinsic stacking fault on

Fig. 12. Photographs (0.5 s exposure) from video monitor showing the gradual disappearance of an extrinsic stacking fault in CdTe [50]. Time intervals of (a) 0 s; (b) 48 s; (c) 85 s; (d) 105 s.

Fig. 13. Appearance (A) and extension (B) of intrinsic stacking faults in ZnTe observed at 500 kV. Time difference of about 3 rain [51].


movement by slip of a Shocldey partial disloca- tion; (iv) diffusional cl imb associated with extrinsic stacking faults.

s imilar defect modif icat ion was also seen in other I I - V I semiconductors ZnSe and ZnTe [51]. The images in fig. 13, for example, were recorded at 500 kV with a time interval of about 3 min, and show bo th the product ion and extension of intrinsic stacking faults in ZnTe due to the slip of Shockley partials.

6. Conclusions

Whilst the realisation of a tomic resolution in the electron microscope is a considerable technical achievement, our experience has shown that skill in operat ing technique, allied with a detailed knowledge of the imaging process, is also essential before useful quanti tat ive informat ion about materials can b e obta ined on a routine basis. Wi th the recent trends t o w a r d s improved (ultra-high) vacuum, in situ specimen t reatment and on-line interactive processing via video camera and computer facilities, these basic H R E M requirements should no t be overlooked. Computer -a ided image simulation and processing are becoming increasingly indispensable for reliable image interpretation.

Acknowledgements

This work has been supported by the Science and Engineering Research Counci l (UK). We are grateful to m a n y friends and colleagues for their valuable collaborations, and we w o u l d part icularly like to thank Dr. V.E. Cosslett and Dr. W.M. Stobbs for their advice, support and encourage- ment throughout this work and Dr. K , C . A . Smi th and Dr. S.J. Erasmus for their involvement with the on-line processing work.

References

[1] R.M. Glaeser, J. Microscopy 117 (1979) 77. [2] V.E. Cosslett, Prof. Phys. Soc. (London) 58 (1946) 443.

[3] Dfl. Smith, R.A. Camps and L,A. Freeman, Inst. Phys. Conf. Sex. 61 (1982) 381.

[4] W.C. Nixon, H. Ahmed, C.J.D. Catto, J.R.A. Cleaver, ICC.A. Smith, A.E, Timbs, P.W. Turner and P.M. Ross, Inst. Phys. Conf. Ser. 36 (1977) b3.

[5] W.C. Nixon, H. Ahmed, C.J.D. Catto, J.R.A. Cleaver, K.C.A. Smith, A.E. Timbs and P.W. Turner, Inst. Phys. Conf. Set. 52 (1980) 443.

[6] V.E. Cosslett, Proc. Roy. Soc. ( ~ d o n ) A370 (1980) 1. [7] DJ. Smith, R.A. Camps. V.E, Cosslett, L.A. Freeman,

W.O. Saxton, W.C. Nixon, H. Ahmed, C.J.D. Catto, J.R.A. Cleaver, K.C.A. Smith and A.E. T'tmbs, Ultramieroscopy 9 (1982) 203.

[8] D.J. Smith, R.A. Camps, L.A~ Freeman, R. Hill, W.C. Nixon and K.C.A. Smith, J. Microscopy 130 (1983) 127.

[9] H. Ahmed, CJ.D, Catto, J.R.A. Cleaver, P,L. Kanitkar and D.J. Smith, J. Phys. E (ScL Instr.) 15 (1982) 1351.

[10] R. Hill and K.C.A. Smith, Inst. Phys. Conf. Set. 61 (1982) 71.

[11] C.J.D. Catto, K.C.A. Smith, W.C. Nixon, S.J. Erasmus and D3. Smith, Inst. Phys. Con£ Set. 61 (1982) 123.

[12] S.J. Erasmus and K.C.A. Smith, J. Microscopy 127 (1982) 185.

[131 D.J. Smith, W.O. Saxton, M~A. O'Keefe, G.J. Wood and W.M. Stobbs, Ultranxieroscopy 11 (1983) 263.

[14] W.O. Saxton, D.J. Smith and S.J. Erasmus, J. Microscopy 130 (1983) 187.

[15] F. Zemlin, Ultramicroscopy 4 (1979) 241. [16] J.M. Cowley and A.F. Moodie, Aeta Cryst. 10 (1957) 609. [17] P.G. Self, M.A. O,Keefe, A.E. Spargo and P.R. Buseck,

Ultramieroscopy 11 (1983) 263. [18] M.A. O;Keefe and P.R. Buseck, Trans. Am. Cryst. Assoc.

15 (1979) 27. [19] A.R. Smith and L. Eyrin8, Ult~an~i'cros~3py 8 (1982) 65. [20] W. Krakow, Mater. Res. Soc. Syrup. Proc. 31 (1984) 39. [21] A.R. Wilson, A.E. Spar8o and D.J. Smith, Optik 61 (1982)

63. [22] W.O. Sexton, and T.L. Koch, J. Microscopy 127 (1982) 69. [23] W.O. Saxton, TJ. Pitt and M. Homer, Ultramieroscopy 4

(1979) 343. [24] J.M. Cook, M.A. O'Keefe, D.J. Smith and W.M. St0bbs, J.

Microscopy 129 (1983) 295. [25] D.J. Smith, L.A. Bursill and GJ. Wood, J. Solid State

Chem. 50 (1983) 51. [26] O. Terasaki, D.J. Smith and G.J. Wood, Aeta. Cryst., in

press. [27] D.J. Smith and M.A. O'Keefe, Aeta Cryst. A39 (1983)

838. [28] M.A. O'Keefe, J.R. Fryer and D.J. Smith, Acta Cryst. A39

(1983) 139. [29] G.J. Wood, W.M. Stobbs and D.J. Smith, "Phil. Mag. A50

(1984) 375. [30] G.J. Wood, unpublished results. [31] S. Iijima and M.A. O'Keefe, J. Microscopy 117 (1979) 347. [32] F. Zemlin, K. Weiss, P. Sehiske, W. Kunath and K.-H.

Hermann, Ultramieroscopy 3 (1978) 49. [33] D.J. Smith, L.A. BursiU and G.J. Wood, Ultramieroscopy

16 (1985) 19.


[34] W.O. Saxton, Inst. Phys. Conf. ~Ser. 36 (1977) 111. [35] J. Frank, Optik 38 (1973) 519. [36] W.O. Saxtun and W. Baumeister, J. Microscopy 127 (1982). [37] M.A. O'Keefe and A. Pitt, in: Proc. 7th European Congr.

on Electron Microscopy, The Hague, 1980 (7th European Regional Congr. on Electron Microscopy Foundation, Leiden, 1980) p. 122.

[38] W.M. Stobbs, G.J. Wood and D.J. Smith, Ultrami- croscopy 14 (1984) 145.

[39] W.O. Saxton and D.J. Smith, Ultramicroscopy 18 (1985) 39.

[40] S. Hovmoller, A. Sjogren, G. Farrams, M. Sundberg and B.-O. Marinder, Nature 311 (1984) 238.

[41] L.D. Marks and D.J. Smith, J. Microscopy 130 (1983) 249. [42] L.D. Marks and D.J. Smith, Nature 303 (1983) 316. [43] D.J. Smith and L.D. Marks, Ultramicroscopy 16 (1985)

101. [44] L.A. Bursill, M.G. Blanchin and D.J. Smith, Proc. Roy.

Soc. (London) A384 (1982) 135.

[45] L.A. Bursill and M.G. Blanchin, J. Physique Lettres 44 (1983) L165.

[46] G.J. Wood, L.A. Bursill and D.J. Smith, J. Microscopy 129 (1983) 263.

[47] L.A. Bursill, Peng Ju Lin and DJ. Smith, J. Solid State Chem. 56 (1985) 203.

[48] D.J. Smith, L.A. Bursill and M.G. Blanchin, Phil. Mag. A50 (1984) 473.

[49] L.A. Bursill and D.J. Smith, Nature 309 (1984) 319. [50] R. Sinclair, F.A. Ponce, T. Yamashita, D.J. Smith, R.A.

Camps, L.A. Freeman, S.J. Erasmus, W.C. Nixon, K.C.A. Smith and C.J.D. Catto, Nature 298 (1982) 127.

[51] R. Sinclair, F.A. Ponce, T. Yamashlta and DJ. Smith, Inst. Phys. Conf. Ser. 67 (1983) 103.

[52] M.A. O'Keefe, in: Proc. 37th Annual EMSA Meeting, San Antonio, TX, 1979, Ed. G.W. Bailey (Claitor's, Baton Rouge, LA, 1979) pp. 556-557.

Documents

Approaching atomic-resolution electron microscopy