Radiation damage and structural studies: Halogenated phthalocyanines

Ultramicroscopy 19 (1986) 279-298 279 North-Holland, Amsterdam

RADIATION DAMAGE AND STRUCTURAL STUDIES: HALOGENATED PHTHAIX)CYAN1NES

David J. SMITH * High Resolution Electron Microscope, Department of Metallurgy, University of Cambridge, Free School Lane, Cambridge CB2 3RQ, UK

J.R. FRYER Department of Chemistry, University of Glasgow, Glasgow G I 2 8QQ~ UK

and

R.A. CAMPS High Resolution Electron Microscope, Department of Metailur~, University of Cambridge, Free School Lane, Cambridge CB2 3RQ, UK

Received 4 February 1986; received in fmai form 14 April 1986

Radiation damage makes it more difficult to accomplish structural studies of beam-sensitive materials. Nevertheless, we report here a 500 kV high-resohition study of halogenated phthalocyanines. Epitaxially grown molecular crystals of fully brominated and fully chlorinated phthalocyanines were found to be isomorphous and had only the a-type polymorph structure. Despite the sensitivity of these crystals to beam damage, the characteristics of thin film growth and details of structurai-defacts could be determined, provided that molecular image resolution was available. The effect of electron beam damage was a degradation of crystallinity which varied considerably according to local morphology and depended on the ease with which the displaced halogen could diffuse away from the parent molecule and thereby avoid recombination. Our f'mdinss demonstrate the feasibility of similar structural studies of other beam-sensitive materials.

1. Introduction

It is well known that covalently-bonded organic molecules are inherently susceptible to radiation damage when irradiated by the high energy flux of an electron microscope and that the halogenated phthalocyanines are unusually radiation resistant. Indeed, using image-averaging techniques, discrete atomic sites within individual molecular columns of the latter have been resolved [1-3]. Although such atomic resolution levels have not yet been obtained in molecular crystals without spatial averaging, details of structural defects and tldn

* Present address: Center for Solid State Science and Depart- merit of Physics, Arizona State University, Tempe, Arizona 85287, USA.

film growth characteristics have been elucidated in several investigations [4-9]. These results were supported by a recent computer-simulation study of molecular crystals [10], where it was shown that molecular positions and orientation should still be recognizable at lower resolution levels ( - 3-4 A), over a range of conditions including specimen thickness and objective lens focus.

The fully halogenated phthalocyanines should be suitable model compounds for investigating whether molecular crystal growth and molecular stacking behaviour can be characterized in other beam-sensitive materials. In this paper, we first review previous electron microscopical studies of halogenated phthalocyanines. We then describe our recent high-resolution observations of fully chlorinated and fully brominated samples, era-

0304-3991/86/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

280 D.J. Smith et al. / Halogenated phthalocyanines

phasizing obvious differences in behaviour, in particular with respect to radiation damage processes. Some of these results have been briefly described in previous conference reports [3,11].

2. Previous work

The pioneering studies of Uyeda and colleagues established the basic structural features of epitaxially grown molecular crystals of Cu-hexade- cachlorophthalocyanine - CuC32NsC116 (CuClPc) [12-14]. It was found, for example, that epitaxy on alkali halides resulted in either thin needles or rhomboidal platelets. The latter crystals contained columns of planar molecules stacked parallel with the substrate, although with lateral offset, as represented schematically in fig. 1. Crystal tilting by about 26 ° gave molecular superposition and, with suitable electron microscope imaging conditions, micrographs were recorded which were directly interpretable as molecular images [12-14] for the first time with organic materials. Observation of the fading of spots in selected area electron diffraction patterns was used as a measure of the loss of molecular integrity due to radiation damage and it was thereby established that the chlorinated compound was 30-40 times more damage-resistant than the non-chlorinated version and that observation at higher electron energies (i.e. higher voltages) substantially reduced the rate of radiation damage, with a factor of about two improve- ment between 100 keV to 500 keV [13]. Careful calibration and measurement of selected area diffraction patterns also enabled the crystal structural data to be determined [14]:

a ffi 19.6.4, b ffi 26.14, c-- 3.68.4, fl -- 116.5°;

space group: C2/c or C2/m.

Subsequent work has mainly been concerned with the characterization of radiation damage and the

discrimination of individual atoms within molecular images. In the former case, Murata and colleagues [15-17] have used optical diffraction patterns, which were taken from high-resolution images of areas as small as 100 A squared, rather than conventional selected area diffraction pat, terns which originate from micrometer-sized regions. Local anisotropy of the damage process could then be established [15-17] and it was found that the degradation of the crystal structure tended to propagate preferentially along defects and in certain crystallographic directions, primarily [100] and [110], accompanied by considerable lattice distortion. Damage processes in CuPc and chlorinated derivatives have also been quantified by digital measurement of electron diffraction pattern spot intensities [18] and the damage mechanisms elucidated by encapsulation techniques [19,20].

The atomic discrimination studies have all been at 500 keV and have principally involved the molecule CuC116Pc [1,3,21-23] although some observations of CuClsBrsPc, described as isomorphous with CuCl16Pc, have also been reported [2]. Since the electron microscopy of these molecular crystals normally includes image recording at typi- cal magnifications of 100-200,000 times, under comparatively low dose imaging conditions (optical density 0.3-0.5), photographic (statistical) noise and support f i l l granularity generally restrict image clarity and make it difficult to dis- tinguish individual atomic column positions [2,11]. With the assistance of photographic averaging (for example, ref. [1]), and/or digital image processing [3,23], greatly improved definition of molecular details can be obtained, in particular with the halogen positions becoming visible. Indeed, a preferential occupation of four specific positions in CuClsBrsPc by bromine atoms has been pro- posed [2] 9 n the basis of variations in Oantrast of the different atomic columns.

Fig. 1. Schematic representation illustrating the stacking of planar molecules parallel to the substrate with lateral offset. Molecular superposition is obtained by tilt of crystal by - 26 °.

3. Experimental

The chlorinated phthalocyanine materials which are available from manufacturers normally con- rain 14-16 chlorine atoms per molecule, thus mak-

D.J. Smith et al. / Halogenated phthalocyanines 281

ing purification necessary before specimen preparation for microscopy. This was done by re-sublimation (twice) in a stream of carbon dioxide and collection of the highest temperature fraction. A KCI single crystal for use as an epitaxial substrate was cleaved on (100) in air and heated to 773 K under vacuum at 10 -6 Torr for 1 h prior to evaporation. The substrate was allowed to cool to 543 K and the CuC1Pc evaporated onto it from a clean Mo boat - the temperature of 543 K had been previously found to give good conditions for epitaxial crystal growth. The calculated average crystal thickness was usually about 100 ,A. After coating with a thin film of carbon, the KC1 was dissolved in distilled water and the carbon/ phthalocyanine film was collected on a fine mesh specimen grid ready for microscopy.

Most observations reported here were made with the Cambridge University high resolution electron microscope [HREM] [24] at an accelerating voltage of 500 kV. At this voltage, there was less electron beam damage than occurs at 100 kV and, because of reduced grain size, photographic emulsion granularity was less restrictive on the minimum useful image magnification. The improved resolution to better than 2 A was also advantageous for resolving submolecular detail in spatially averaged micrographs. The microscope was equipped with a + 30 ° double-tilt specimen holder and a low inertia beam shutter device above the sample which facilitated the minimum exposure techniques used extensively throughout this work [25], and an on-line image pick-up and viewing system [26] was particularly useful for preliminary focussing at very low beam currents. Electron doses impinging on the sample were measured either with the beam shutter, which could also be used as a Faraday cage, or with the final viewing screen, though with less accuracy. Images were usually recorded on Agfa Scientia 23D56 EM film developed under standard conditions. A JEM 100C HREM was used for preliminary observations and for characterizing the changes in morphology caused by deposition and annealing at different substrate temperatures.

4. Results

4.1. Thin film growth and annealing

In preparation for the study reported here, preliminary investigations were first made of the influence of different substrate temperatures, and subsequent annealing, on the quality of the thin molecular films grown by epitaxy [27]. It was found, for example, that there was no visible structure in the apparently amorphous phthalocyanine films grown with the substrate held at room temperature whereas, for a temperature of 459 K, some ordering was starting to occur. At higher temperatures, larger crystals were formed which, because of their epitaxial growth, invaria- bly had their principal axes aligned rather than randomly orientated. They were generally found to be wedge-shaped, having somewhat irregular contours at about 543 K but developing into more rectangular profiles at 618 K. Smaller rod-shaped crystals were also common. These did not appear to have any preferred crystal growth habit relative to the substrate normal although they were again usually aligned along either of two orthogonal axes in the plane of the substrate, as shown in the low magnification image of CuBrPc in fig. 2, thus indicating that they were influenced to some ex- tent by the substrate. The appearance of low-contrast tracks across the substrate leading to these crystals in fig. 2a suggests that they might have been mobile. However, imaging at high magnification, as shown in fig. 2b, indicates that these "tracks" correspond to long and thin needle- shaped crystals. In the extended platelet or spatulate crystals where widespread epitaxial growth had occurred, the usual procedure for locating the molecular superposition axis was to monitor a slightly-defocussed selected area diffraction pattern (SADP) as the crystal orientation relative to the incident beam direction was slowly altered. Close to the zone axis orientation the contrast of a crystal became very dark and a characteristic, densely populated, array of spots was seen in the SADP, as shown in figs. 3a and 3b which are for CuBrPc and CuC1Pc respectively. These patterns confirm that the two materials are isomorphous although measurement indicates slight differences


Fig. 2. (a) Low magnification image of Cu-hexadecabromophthalocyanine on thin carbon support film showing a large extended crystal and other smaller, basically rod-shaped, crystals. Note the narrow, almost orthogonal, low contrast lines and depletion of substrate around large crystal. (b) Higher magnification image of CuBrPc showing small crystalfites, as well as long and narrow sets of lattice fringes arising from needle-shaped crystals.

in unit-cell dimensions. It is significant that no POlymorphs of either the chlorinated or brominated materials have been observed.

4.2. High-resolution structure

Simulated HREM images of CuC1Pc have previously established that the position and orienta-

Fig. 3. (a) Selected area electron diffraction pattern from CuClPc. (b) Selected area electron diffraction pattern from CuBrPc.

tion of molecular columns should be unambiguous at an accelerating voltage of 500 kV for crystal thicknesses up to about 200 A, even though sec- ond-order contributions are then causing considerable artefactual image detail (see fig. 6a of ref. [10]). This was confirmed in experimental micrographs, such as the through-focal series of images from a perfect crystalline region shown in fig. 4. Furthermore, it is significant that, despite rever- sals in contrast of the detail in these images, the "dover-leaf" or quatrefoil shape of individual molecules is generally recognizable throughout the series. This alleviates somewhat the usual require- ment of high-resolution imaging that micrographs should be recorded close to the optimum defocus (i.e. over a range from ~ -500 to -900 .~) which would otherwise substantially restrict the usable area of the micrograph. (It should be realized that, because of the specimen tilt of 26.5 ° required for molecular superposition, there could be a focal difference in excess of 5000 & across the specimen area recorded at normal magnifications,)


Fig. 4. Through-focal series (800 A focal steps) of 500 kV high-resolution images of CuClPc. Positions of molecular columns recognizable in each image.

Discrimination of individual atomic columns within separate molecular columns is normally restricted by the poor signal-to-noise ratios associated with the limited beam current required for low-dose imaging. This is apparent in the regions of CuBrPc and CuC1Pc, shown at very high magnification in figs. 5 and 6a respectively, which were recorded close to the optimum defocus at initial magnifications of - 120,000 times. Follow- ing some local signal averaging and filtering, as

shown in fig. 6b, much improved image definition is obtained and, in particular, every chlorine atom position around the periphery of the molecule can be easily recognized. Fig. 6c shows, in contour form, the atomic and molecular image which results after 4-fold (nun) symmetry is imposed on the filtered micrograph.

It was interesting that small fringe "packets" were observed in some areas. Although these packets often appeared as a result of radiation


2 ' ~ : : : : . . . . . : +

from the crystal growth process. Electron diffraction confirmed this identity, with Cu20 and CuO both being present [28].

4.3. Structural defects

Fig. 5. High magnification image of CuBrPc recorded close to optimum defocus.

damage processes, a phenomenon which will be described in section 4.4 below, the fringes were also found in regions, such as that shown in fig. 7, which had not previously been exposed to the electron beam. With lattice spacings typically of

o

about 2.5 A, the packets are believed to be small impurity crystallites of a copper oxide, arising

Whilst image averaging is obviously not appli- cable to aperiodic specimen features such as dislo- cations or interfaces, it is still to be expected, given the comparative insensitivity of molecular images to focus and thickness changes, that useful morphological information on the molecular scale should still be obtainable. Fig. 8a, for example, shows the simplest type of stacking fault which results when the interface between two adjoining phthalocyanine crystals, which have merged together, lies exactly parallel to a <110) crystallographic plane but with a lattice offset remaining at the junction. In such cases, it is evident from the image that both crystals remain essentially perfect except in the immediate + vicinity of the interface where the mismatch appears to affect a strip of, at most, two molecular diameters in width. Within this strip, it is common for the image contrast to be lower than it is in the two adjacent crystals and there is often some molecular rotation, as seen at the interfaces shown in figs, 8b and 8<:. Note also the thin white fringe (an'owed) in these latter two images, lying along the interface, which is probably a Fresnel-like effect caused by a slight gap between molecular columns at the junction. It is significant that even this simple interface occurred only infrequently in the epitaxial Crystals of the brominated sample, despite lateral crystal dimen-

Fig. 6. (a) High magnification image of CuC1Pc after image digitization. (b) Image enhancement as a result of local signal averaging and filtering. (c) Resulting atomic and molecular image in contour form after imposing 4-fold symmetry on processed image. Chlorine atomic column positions dearly visible.


Fig. 7. Region of CuCIPc showing small packet of lattice fringes (circled).

sions typically of more than 0.5/~m. More complicated boundaries between crystals

were commonly found in the chlorinated sample, with the complexity seeming to depend on the mean orientation of the boundary relative to ~110) direction. Fig. 9a shows an interface having a single sidestep, one molecular column wide, with a region of local disorder at the position where the actual sidestep occurs. Molecular rotation and

brighter line contrast are again apparent along the two separate sections of the interface. Further boundaries containing sidesteps of 2 and 4 molecular columns, respectively, are shown in figs. 9b and 9c. These two interfaces differ from that shown in fig. 9a in that the crystal structure is locally different, although remaining crystalline, at the points where the actual sidestep takes place. Higher magnification images indicate that the interfaces appear to contain extra half-molecules at most of these positions. A .schematic representation of the molecule configuration at these sidesteps and other boundaries is described later (see section 5.1).

The grain boundary shown in fig. 10a sidesteps by a total of 9 molecular columns over its length at the positions marked A to F and there is considerable variation in the way that these sidesteps appear to be accommodated. In the enlargement shown in fig. 10b, which is from the region of the interface labelled A, B, C, the quatrefoil molecular shape is not visible along the boundary. Instead, it appears that the tilt of the molecules is altered relative to those in the neighbouring bulk crystal and it is not immediately clear from the image where, or even how, the sidesteps near A and B occur. The double sidestep at E, whilst more localized, also features these apparently inclined molecules as well as extra half-planes as seen previously in figs. 9b and 9c. The triple sidestep at D corresponds to a region where the average local orientation of the boundary is almost around to [010]. The positions of the molecu-

Fig. 8. (a) Interface between two CuCIPc crystals with lattice off-set. Note slight rotation of molecular columns along the interface. (b), (c) As for (a) but note the strip of brighter contrast along the interface.


Fig. 9. Interfaces between adjacent crystals of CuC1Pc: (a) with single sidestep and some amorphous material; (b) sidestep by two molecular columns in width; (c) sidestep by four molecular columns in width.

lar columns are not distinct and a patch of finely spaced lattice fringes becomes visible in later images of this area suggesting that susceptibility to radiation damage is enhanced in this region where there is considerable molecular disorder due to the severe mismatch across the boundary.

One of the few grain boundaries seen in the fully brominated sample is shown in fig. 11; it was located near one edge of a crystal which overall was more than 0.6 x 0.4 #m 2 in area. In the region of A and B, shown at higher magnification in fig. 11b, there are two short segments having an interface lying parallel to [110] where the (110) fringes are offset. Between A and B the lattice is dis- continuous and there is a small patch (arrowed) which seems amorphous. It is also interesting that at C the crystal itself is free of defects and yet there is a continuous lattice around both sides of the amorphous region at D extending from the interface at B, which implies that there must be some lattice distortion between B and C.

It was not uncommon to find similar interfaces in the chlorinated sample, having a good lattice match on one side of a hole but with an offset on

the other, although the means by which the associated lattice distortion took place seemed variable. In fig. 12a, for example, sighting along the direction of the arrow reveals that the (110) fringes at the top right are basically rotated by several de- grees. In fig. 12b, the lattice simply bows outwards on the right-hand side. Also note here, by viewing along the arrow, the extra molecular column which terminates at A, which thus corresponds to the classic configuration for an edge dislocation, viewed end-on.

4. 4. Damage processes

The recent work of Fryer and Holland [19] established that encapsulation generally reduced the overall susceptibility of molecular crystals to electron-beam damage. In this section, however, our primary objective is to consider the influence of local specimen morphology on radiation semi- tivity. It is also of interest to establish whether observations at higher accelerating voltage result in any substantive differences compared with the earlier studies of radiation damage processes in


i

Fig. 10. (a) Compficated grain boundary ABCDEFG in CuC1Pc which sidesteps by a total of nine molecular columns over its length. (b) Enlargement from (a) showing region ABCDE at higher magnification.

CuC1Pc which were all carried out at an accelerating voltage of 100 kV [15-17].

Figs. 13a-13c show three successive micrographs of a CuC1Pc crystal recorded at the same objective lens defocus under minimum exposure conditions, i.e. with the beam incident on the sample only while the image was being taken. The respective cumulative electron dose of each image was 0.8, 1.7 and 2.6 C/cm 2. In the bulk of the crystal there is no noticeable difference in the appearance of the molecular images. However,

along the outside edges, and around the amorphous regions within the crystal, there is considerable loss of crystallinity, which can be clearly seen by comparing the lattice planes arrowed at A, B and C in each image and the reference point X. It also appears that the loss of molecular definition at D is isotropic, i.e. without any preferred orientation.

The series of images in figs. 14a-14c show a grain boundary, similar to those in section 4.3, which has an offset of the molecular columns across the interface, as well as sidesteps (arrowed)


Fig. 11. (a) Region of CuBrPc showing two short interfaces with lattice offsets at A and B but with no offset at C. (b) Higher magnification image of (a).

at A and B. (The lack of mm symmetry here is due to local crystal tilting off the zone axis.) These images were taken over a longer time interval than those in fig. 13, with the shutter above the sample left open between exposures, and the cumulative doses here are 1.0, 4.4 and 7.6 C/cm 2 respectively. It is again apparent that the crystallinity at the edges degrades much faster than does the bulk, whilst the interface itself has an intermediate rate of damage, although damaging more rapidly at the

sidestep at A than along the straight section between A and B. Finally, despite the damage at the crystal edges and interfaces, the molecular images in the bulk remain well-defined, though slightly asymmetric due to local crystal tilt.

As mentioned previously, the radiation damage observed in the chlorinated material was accompanied by the occurrence of copper oxides, in the form of small crystals with distinct lattice fringes. These crystals were found to retain their crystal-

D.J. Smith et al. / Haiogenated phthalocyanines 289

Fig. 12. (a) Interface in CuCIPc with good lattice match on one side of amorphous (depleted) region but with offset at the other; note slight apparent local rotation of lattice (sight along arrow). (b) As for (a) but note the extra row of molecular columns (arrowed) which terminates at A.

linity even after the surrounding phthalocyanine crystal had been amorphised. The initial development of one such area is shown in the pair of images in fig. 15. (This region is in reverse contrast, with white molecular columns, due to an over-focus lens current condition.) In fig. 15a, the crystal has somewhat lower contrast and slightly disordered molecular columns at the position marked X which have been replaced, in fig. 15b, by an approximately rectangular patch of fringes about 30 x 50 /~2 in size. The pair of images in

figs. 16a and 16b, here in "correct" contrast, show a similar-sized area of fringes ( m o w e d ) after the sample has received electron doses of 1.0 and 8.2 C / c m 2 respectively. It is interesting that, in fig. 16b, the disintegration of the CuC1Pc crystal has

• clearly worked inwards from the edge of the crystal around, and past, this crystaUite. During this process, the size of the crystallite does not, however, appear to have altered.

There were differences and similarities with the brominated and chlorinated samples with respect


Fig. 13. Successive images of CuC1Pc recorded at fixed objectives lens defocus under mimmal exposure conditions. Cumulative electron doses of 0.8, 1.7 and 2.6 C / c m 2. Note the general loss of crystallinity around the edges and holes which is made most apparent by viewing along the arrows at A, B and C.

D.J. Smith et al. / Haiogenated phthalocyanines 291

Fig. 14. Successive images of a grain boundary in CuCIPc which has both offsets and sidesteps. Cumulative electron doses of 1.0, 4.4 and 7.6 C/era 2. Note that the differences in rates of electron beam damage depend on local environment (see text).

to the electron beam damage processes. Copper oxide crystals as observed in the un-irradiated CuC1Pc sample (fig. 7) were not seen, nor did they develop following irradiation (fig 15). On the other hand, the preferential destruction along crystal edges and around small holes, rather than in the crystal bulk or along interfaces, was again apparent. The three images shown in figs. 17a-17c, for example, are from the brominated sample and have been recorded with identical doses to those of the chlorinated sample shown in figs. 14a-14c, i.e. 1.0, 4.4 and 7.6 C/cm 2 respectively. Although these images come from a region of the sample which is well underfocus, the molecular shape and

orientation can still be clearly recognized, as can the offset across the interface (arrowed). The ex- tent of crystal disintegration seems more developed in fig. 17c than in the comparable fig. 14c. However, it should be appreciated, after table II of ref. [19], that electron doses for extinction of diffraction pattern spots can differ by factors of 3 or more with only a doubling of crystal thickness. Moreover, for chlorinated and brominated crystals known to be of almost equal thickness, extinction doses were found to be almost the same. Hence, we do not believe the differences here between samples to be significant.


5. Discussion

5.1. Crystal growth mechanisms and structural defects

Fig. 15. Development of small copper oxide crystal in CuCIPc as a result of electron beam damage (at reverse contrast focus position).

There are essentially two different ways in which effectively planar molecular crystals can stack together and these polymorphs are illustrated in fig. 18a. In the a-phase (left), all the molecules are stacked in parallel, whereas in the ~-phase (right), they have a herring-bone configuration. Other slight variants, also known as polymorphs, have been reported in metal phthalocyanines (for example, refs. [28,5]), but we have only obtained evi- dence for the a-phase in the fully halogenated materials observed during this study. This is presumably because the presence of the halogens increases the anisotropy of the intermolecular forces thereby making the occurrence of other polymorphs energetically unfavourable.

The epitaxial growth of molecular crystals on a substrate can also happen in a number of ways - the possibilities for the a-phase polymorph are represented in fig. 18b - and the actual mode adopted will depend on a number of factors including the nature of the substrate, in particular its temperature and cleanliness, the strength of the intermolecular forces a n d the interaction of the molecules with the substrate. Clearly, the preparation of a substrate as a suitable surface for the growth of large crystals free of defects requires that it contain few nucleation sites and this, in turn, demands a compromise between providing sufficient heat to degas impurities and yet avoid- ing excess sublimation of the substrate. From fig. 2a, it would appear that the pre-treatment used for cleaning the KC1 substrate has in fact evaporated some of it, leaving square or rectangular domains, and the needle-shaped phthalocyanine crystals have aligned along the resulting terrace ledges. In flatter areas there should be less hindrance to molecular mobility across the surface, thus permitting the formation of the spatulate

Fig. 16. Small copper oxide crystal which remains intact even after substantial dectron beam damage to surrounding phthalocyanine crystal.

D.J. Smith et aL / Halogenated phthalocyanines 293

O

Fig. 17. Development of electron beam damage in CuBrPc. Cumulative electron doses of 1.0, 4.4 and 7.6 C/cm 2. Compare with corresponding part of fig. 14 and see text.

crystals, and this is also consistent with the observation that the substrate was generally bare of the small needle growths in these regions. It is also relevant here to mention again the observation of section 4.1 that the generally random orientation of crystals prevalent for an evaporation onto a substrate at 459 K did not occur for higher tem-

perature anneals. This indicates that, in the former case, there is still insufficient interaction with the substrate to result in crystal rearrangement, even though surface diffusion has enabled crystal growth to take place. At room temperature, there was not enough energy even for surface diffusion.

At the substrate temperature eventually used


, ~ r 7,,i¢',~"

\ \ \ \ \ \ \ \ ' , , ~ \ \ \ \ ~ \ \ \ \ \ \

b

Fig. 18. (a) Schematic representation of a- and /]-polymorph structures. (b) Schematic showing possible variants in stacking of a-polymorph.

for epitaxy, namely 543K, the extended crystals grew easily and were well aligned along orthogonal axes, presumably corresponding to the ~100) axes of the substrate. The characteristics of crystal growth are then effectively dictated by the shape of the phthalocyanine molecules and their aniso- tropic intermolecular forces. As crystals grow outwards from their nucleation sites on the substrate they will eventually merge together and it might then be anticipated that an offset of lattices at the junctions will frequently occur, as was in fact observed. A simple offset would result in an energetically unfavourable location of molecules, particularly with respect to the halogen positions, as well as considerable strain, and the experimental micrographs of fig. 9 suggest that this situation is usually relieved by a slight rotation of the molecules and a change in the angle of tilt of the molecular stacks in the two rows of molecular columns immediately adjacent to the boundary relative to those in the bulk crystal. This is represented schematically in planar view in fig. 19. It is significant, however, that no lattice offsets were observed which would have corresponded to the boundary being a non-integral number of molecular columns in width, despite a high probability for this to occur. There would obviously be considerable strain associated with this configuration but presumably the epitaxial temperature provides sufficient molecular mobility for these defects to be annealed out.

More complicated molecular configurations oc-

XxxX

Fig. 19. Schematic planar view representing molecular column stacking at an interface having offset only.

cur when the junction between merging crystals is not conveniently aligned along ~110) and in some instances it is clearly impractical to deduce the detailed nature of the molecular stacking accom- panying such interfaces, The interface A - G shown in fig. 10 effectively demonstrates the limit to which this can be done. Between E and F, for example, a s well as between F and G, this boundary resembles the simple coherent interfaces seen in fig. 9 although with a side-step, i.e., lateral displacement, at F apparently accompanied by an extra half molecular image. This sidestep can be visualized as resulting from a molecular column which is out of register by about a half molecule with the molecular columns above it. This is represented schematically in fig. 20, in both planar and cross-sectional projections, and corresponds to an edge dislocation, one molecule in width, having a Burgers vector lying along the boundary. The planar view of the molecular stacks shows clearly how the offset in the column positions across the interface is altered by the presence of the dislocation.

At the interface near A, B and C, careful measurement of the locations of the molecular columns, by reference to the positions of the chlorine atoms which are the major source of image contrast, indicates that the columns must be inclined at about an extra 12 ° relative to those in the surrounding bulk crystal (which is already tilted by -26½ ° in order to get molecular superposition). This reduced area of molecular projection also means that there is increased space between the molecules along the boundary which is esti- mated to be about 3 A per molecule. The sidestep


Fig. 20. Schematic in both planar and cross-sectional view of an interface having both offsets and sidesteps as, for example, visible in fig. 10.

at C can again be represented by an edge dislocation, similar to that shown in fig. 20, with the difference that the molecules at the boundary have increased tilt which would result in their protrud- ing well above the surface of the crystal thereby making them more susceptible to chemical attack.

5.2. Mechanisms for radiation damage

Most studies of radiation damage in organic materials have been concerned with the decay of diffracted beam intensities as a function of electron dose. Whilst the sensitivities of different materials can, in principle, be established by comparing the relative rates of loss of their crystallinity under carefully controlled experimental conditions, localized information about the nature of the radiation damage process, nevertheless, remains strictly limited, or even unobtainable, using this technique. This constraint obviously does not hold for high resolution imaging on the molecular scale although similar restrictions apply to other bulk characterization methods such as mass-loss

spectrometry, or energy-dispersive X-ray micro- analysis, despite the useful insights which can sometimes be provided. For example, it was shown by the latter technique that at a stage when all diffraction spots from a specimen of CuClltPc had disappeared, thereby indicating no remaining crystallinity, there had in fact been a loss of - 15% of the chlorine atoms from the bulk of the sample and a substantial proportion of the remainder were presumed to have diffused throughout the sample [17]. Furthermore, the detailed diffraction study of molecular crystals [18] has shown that encapsulation of the specimen between carbon and other thin films leads to a significant decrease, typically by factors of 3 or more, in rates of radiation damage relative to uncoated samples. These authors propose that the effect of the protective coating is to restrict the diffusion of atoms away from the damaged molecules and thereby enhance the probability of recombination taking place. This suggestion is also consistent with earlier low magnification imaging of two chlorinated derivatives of CuPc where it was observed [17] that the visible effects of radiation damage propagated into the bulk from crystal edges and holes.

The development of radiation damage seen in our experimental micrographs (see section 4.4) can also be rationalized in terms of an encapsulation or "cage" effect. Thus, in fig. 13, the high-resolution detail first starts to disappear around the crystal edges, presumably because of the increased probability of any displaced chlorine atoms there being able to diffuse out of the crystal. In the crystal bulk, the molecules are more tighly bound and, since this makes it more difficult for the large anions to move around, recombination is much more likely. At crystal defects, molecules are not generally as well packed and one would therefore predict some sort of intermediary state with re- gard to damage rates and this was, in fact, what was observed. For the grain boundaries shown in figs. 14 and 16, for example, it was in the regions of the sidesteps where there was considerable lattice strain, i.e. greatest molecular mismatch, that degradation developed most rapidly. Once the damage has been initiated at such defects, they will act as "damage centers" where diffusion is


enhanced and the damage in effect becomes self- propagating.

Finally, it is worth noting that this large variation in damage rates according to local defect structure, as well as its strong dependence on thickness [19], probably explains much of the considerable confusion which can result from comparing published critical dose studies. This point is discussed in detail elsewhere [20].

6. Conclusion

Our observations of fully halogenated phthalocyanines have established that they have significant differences relative to other molecular crystals. Unlike metal- and metal-free-phthalocyanines, for example, epitaxial growth does not result in a variety of polymorphs and many common growth defects such as vacancy loops and twin boundaries are not observed (see, for com- parison, refs. [6-7,30]). Mismatch occurs at grain boundaries but, unlike quaterylene [9], it seems to be accommodated only by molecular column rotation and tilt rather than by bending and distortion of individual molecules. These structural characteristics are obviously related to the strong influence of the halogens on the anisotropy of the intermolecular forces which thereby make other stacking configurations energetically unfavourable. In effect, the presence and size of the halogens substantially reduce the options available for stacking the molecules together.

Finally, it can be concluded that it should be possible to undertake similar high-resolution structural studies of other radiation-sensitive materials. Ultimate success will, however, depend upon the particular damage mechanism(s) and upon whether any means can be found for reduc- ing the damage rate, e.g. encapsulation, lower temperature, etc.

Acknowledgements

We are grateful to the Science and Engineering Research Council for provision of the JEM 100C at Glasgow, for support of the Cambridge Univer-

sity 600 kV high resolution electron microscope project, and for support of commuting between our establishments. We also thank Dr. W.O. Sax- ton for image processing.

References

[1] N. Uyeda, T. Kobayashl, K. Ishizuka and Y. Fujiyoshi, Chem. Scripta 14 (1978-79) 47.

[2] T. Kobayashi, Y. Fujiyoshi, K. Ishizuka and N. Uyeda, in: Electron Microscopy 1980 (Proc. 7th European Congr. on Electron Microscopy, The Hague, 1980), Vol. 4, Eds. J. Van Landuyt and (3. Boom (7th European Congr. on EM Foundation, 1980) p. 158.

[3] J.R. Fryer, D.J. Smith and R.A. Camps, in: Electron Microscopy and Analysis 1983 (Proc. Intern. Conf., Guildford, 1983), Inst. Phys. Conf. Ser. 68, Ed. P. Doig (Institute of Physics, London-Bristol, 1984) p. 15.

[4] Y. Murata, T. Baird and J.R. Fryer, Nature 262 (1976) 721.

[5] T. Kobayashi, Y. Fujiyoshi, F. Iwatsu and N. Uyeda, Acta Cryst. A37 (1981) 692.

[6] T. Kobayashi, Y. Fujiyoshi and N. Uyeda, Acta Cryst. A38 (1982) 356.

[7] T. Kobayashi, Y. Fujiyoshi and N. Uyeda, J. Crystal Growth 65 (1983) 511.

[8] J.R. Fryer, Acta Cryst. A35 (1979) 327. [9] J.R. Fryer and D.J. Smith, Proc. Roy. Soc. (London) A381

(1982) 225. [10] M.A. O'Keefe, J.R. Fryer and D.J. Smith, Acta Cryst. A39

(1983) 838. [11] J.R. Fryer, R.A. Camps and D.J. Smith, in: Electron

Microscopy 1982. (Proc. 10th Intern. Congr. on Electron Microscopy, Hamburg, 1982) (Deut. Ges. fl~r Electro- nenmikroskopie, 1982) Vol. 2, p. 449.

[12] N. Uyeda, T. Kobayashi, E. Suito, Y. Harada and M. Watanabe, in: Microscopie Electronique 1970 (Proc. 7th Intern. Congr. on Electron Microscopy, Grenoble, 1970), Ed. P. Favard (Sot. Fran~fise de Microscopie Electro- nique, Paris, 1970) Vol. 1, p. 23.

[13] N. Uyeda, T. Kobayashi, M. Ohara, M. Watanabe, T. Taoka and Y. Harada, in: Electron Microscopy 1972 (Proc. 5th European Congr., Manchester, 1972), Inst. Phys. Conf. Set. 14, Ed. A.M. Glanert (Institute of Physics, London-Bristol, 1972) p. 566.

[14] N. Uyeda, T. Kobayashi, E. Suito, Y. Harada and M. Watanabe, J. Appl. Phys. 43 (1972) 5181.

[15] Y. Murata, T. Baird and J.R. Fryer, Nature 262 (1976) 721.

[16] Y. Murata, J.R. Fryer, T. Baird and H. Murata, Acta Cryst. A33 (1977) 198.

[17] Y. Murata, in: Electron Microscopy 1978 (Proc. 9th In- tern. Congr. on Electron Microscopy, Toronto, 1978), Ed. J.H. Sturgess (Microsc. Soc. of Canada, 1978) Vol. 3, p. 49.


[18] W.R.K. Clark, J.N. Chapman, A.M. MacLeod and R.P. Ferrier, Ultramicroscopy 5 (1980) 195.

[19] J.R. Fryer and F. Holland, Proc. Roy. Soc. (London) A393 (1984) 353.

[20] J.R. Fryer, Ultramicroscopy 14 (1984) 227. [21] N. Uyeda, E. Kirkland and Y. Fujiyoshi, in: Electron

Microscopy 1978 (Proc. 9th Intern. Congr. on Electron Microscopy, Toronto, 1978), Ed. J.H. Sturgess (Microsc. Soc. of Canada, 1978) Vol. 1, p. 242.

[22] N. Uyeda, E. Kirldand, Y. Fujiyoshi and B.M. Siegel, in: Electron Microscopy 1978 (Prec. 9th Intern. Congr. on Electron Microscopy, Toronto, 1976), Ed. J.H. Sturgess (Microsc. Soc. of Canada, 1978) Vol. 1, p. 220.

[23] E.J. Kirkland, Ultramicroscopy 9 (1982) 45. [24] D.J. 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. Timbs, Ultramicroscopy 9 (1982) 203.

[25] J.R. Fryer, J.R.A. Cleaver and D.J. Smith, in: Electron

Microscopy and Analysis 1979 (Proc. Intern. Conf., Brigh- ton, 1979), Inst. Phys. ConL Ser. 52, Ed. T. Mulvey (Institute of Physics, London-Bristol, 1980) p. 207

[26] C.J.D. Catto, K.C.A. Smith, W.C. Nixon, S.J. Erasmus and D.J. Smith, in: Electron Microscopy and Analysis 1981 (Proc. Intern. Conf., Cambridge, 1981), Inst. Phys. Conf. Set. 61, Ed. M.J. Goringe (Institute of Physics, London-Brighton, 1982) p. 123.

[27] J.R. Fryer, in: Proc. 37th Annual EMSA Meeting, San Antonio, TX, 1979, Ed. G.W. Bailey (Claitor's, Baton Rouge, LA, 1979) p. 616.

[28] F. Holland, PhD Thesis, University of Glasgow (1984). [29] J.R. Fryer, in: Developments in Electron Microscopy and

Analysis 1977 (Proc. Intern. Conf., Glasgow, 1977), Inst. Phys. Conf. Ser. 36, Ed. D.L. Misell (Institute of Physics, London-Brighton, 1977) p. 423.

[30] Y. Fujiyoshi, T. Kobayashi, K. Ishizuka, N. Uyeda, Y. Ishida and Y. Harada, Ultramicroscopy 5 (1980) 459.

Documents

Radiation damage and structural studies: Halogenated phthalocyanines