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Analytical electron microscopy of thin films and interfaces
154
Analytical electron microscopy of thin films and interfaces R. H. Geiss
Abstract In this paper a brief overview is presented of some of the techniques available with transmission electron microscopy (TEM) to characterize thin films and interfaces. The critical role that sample preparation plays is discussed. Examples of the application of many of the analytical techniques used in TEM are given. These include nano area electron diffraction, energy-dispersive X-ray spectroscopy and electron energy loss spectroscopy. High resolution imaging is available at atomic resolution and new methods are being developed to analyze the images including computer simulation. Imaging of magnetic domains is important to the study of materials used in digital data storage.
1. Introduction
One of the most powerful techniques which can be used in the study of thin films and interfaces is high resolution transmission electron microscopy (HRTEM). When combined with analytical tools such as nano diffraction, energy loss spectroscopy (energydispersive spectroscopy (EDS)) and electron energy loss spectroscopy (EELS), it is given the name analytical electron microscopy (AEM). One of the major hurdles which must be overcome in any AEM investigation is the preparation of a thin sample which is representative of the material under investigation. It is appropriate that a brief discussion of the newer approaches to sample preparation be included. A few of the ways that AEM can be used to study thin films will be explored. Both single-layer (SL) and multilayer (ML) films in the plan and cross-section views will be considered. With modern AEM it is possible to obtain elemental data using EDS and EELS with a lateral resolution of 1- 10 nm depending on the microscope and the sample. Electron diffraction can be used to determine phase information, crystallographic orientation and lattice parameters from areas as small as l-2
nm in diameter
in microscopes
with
nano
probe
can be used to obtain images with better than 0.2 nm point-to-point resolution. To aid in the interpretation of these images computer simulation of the structure at the atomic scale should be included. Lorentz imaging of magnetic domains is very valuable in the study of digital storage media. To illustrate the above, examples will be given which show the use of AEM to determine (a) the grain size and details of the structure with high resolution capabilities.
HRTEM
imaging, (b) the structure of nano size regions in crystalline films with electron diffraction and (c) the use of EDS to determine the thickness of very thin films and also the minimum detection capabilities of EDS. Reference to the use of parallel EELS (PEELS) to determine the oxidation state of thin film oxides will be discussed. The use of Lorentz microscopy to image magnetic domain structures will be shown. Finally, some of the approaches used in the analysis of HRTEM lattice images of interface and defect structures will be discussed. The use of Fast Fourier Transforms (FFTs) and moire patterns may provide a mechanism for making measurements from these images.
2. Sample preparation The technique of specimen preparation of bulk materials has evolved from the methods of chemical and/or electrochemical thinning to techniques which include mechanical slicing, polishing and dimpling and a final thinning using argon ions. The latter methods have the advantage that they are more readily used on samples containing a combination of materials. It is almost impossible to find a chemical which would uniformly thin a sample consisting of an arbitrary mixture of components. The new techniques are also more easily used to prepare samples with localized defects or other important features. Mechanical polishing-ion thinning is the only method which can be used to prepare cross-sectional views reliably from samples on silicon or sapphire substrates. Although the development of these techniques has been particularly valuable in the study of semiconductor structures in cross-section, they can be used for a much broader variety of materials. The
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R. H. Geiss 1 Analytical
reader is referred to refs. 1 and 2 for a thorough discussion of these techniques. Two problems cause major concern, however; specimen heating which occurs during the many stages of the preparation, especially in the curing of the M Bond610, the epoxy recommended for silicon [ 11. The curing procedure for this epoxy suggests heating to 100 “C for 10 h or more. Heating may also occur during the ion thinning, especially if the thinning time is excessive, e.g. more than 12 h. There is also the possibility that during ion thinning some of the material which is removed may be redeposited onto other parts of the specimen. A technique in which samples are thinned by mechanical polishing only or with at most a few minutes of ion thinning [3] may help to alleviate the latter problem, but one must always be very aware of the potential problems associated with prolonged heating. For the preparation of plan section specimens any of the techniques discussed above can be used. However, if the effect of the substrate is not deemed critical, or if the heating is a problem, plan view samples of thin films can be prepared very quickly using a substrate of evaporated carbon on freshly cleaved mica. Very large area samples are obtained by releasing the film at the carbon-mica interface in water. If the exposure to water is undesirable, special thin window substrates such as silicon nitride on a silicon wafer can be used. Here the film is deposited on the Si,N, and used directly for TEM. The difficulty here is in manufacturing suitably thin silicon nitride windows. Films can be also prepared either by direct evaporation onto NaCl or by evaporation onto a substrate film previously deposited on NaCl. However, these films are also retrieved by floating them off in water. In principle, almost any sample can thinned for TEM study given the variety of techniques available today. However, one must be aware that sample preparation may be the most tedious and frustrating part of the study, especially in view of the criterion discussed above.
electron
microscop)
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In a study of the structure of thin evaporated films of copper on carbon substrates [4] we found that films 1, 2 and 3 nm thick all had island-like structures consisting of small crystallites, 2-5 nm in dimension, surrounded by amorphous regions thought to be the carbon substrate (Fig. 1). However, an investigation of the 2 nm film with EDS using a 5 nm probe showed that the amorphous areas between the islands also contained copper. The thickness of the copper at the islands and at the amorphous areas between, as determined by the integrated intensity of the Cu Ka peak, varied from about 1 nm between the islands to more than 4 nm at the islands. The film thickness monitor and electron microprobe measurements had indicated that the average film thickness was right at the 2 nm mark. Even though one might argue the accuracy of the 1 and 4 nm thickness determinations, there is no doubt that the copper intensity on the islands was much greater than that from a large area sample and that there was copper between the islands. These measurements indicate a film profile which is quite different from that expected. Microdiffraction from areas between the islands showed that the structure was noncrystalline, suggesting that the regions between the islands may be amorphous. In an investigation using EDS to determine the minimum detectable limit (MDL) in thin films and alloys, Geiss and Savoy [5] showed that by using the total area under the peak, as the peak intensity, and the intensity of a single channel, as the background intensity, an MDL of about 150 ppm can be measured in most metal systems. A critical step in the analysis is the ability to extract a very small peak from a region of high background in the EDS spectra. This was done using a
3. Analysis 3.1. Single -layer films SL films may be single-element, compound, mixed element or alloy films. In the case of single-element films one is probably concerned with some aspect of the growth mechanism which might include the influence of the substrate, or with defects such as vacancies, dislocations or grain boundaries. In alloy SL films one can also study the structure and properties of different phases as well as the defects etc. To illustrate the use of AEM in the analysis of SL films consider the following examples.
Fig. 1. High resolution electron micrograph showing the island structure of a 2 nm thick copper film on a carbon substrate. EDS showed that the regions between the crystalline islands had as much as 1 nm of copper and that the islands were as thick as 4 nm.
personal computer and importing the EDS data in ASCII format into a spreadsheet template developed by one of these researchers [6]. Patterson et ul. [7] used PEELS to reveal differences in the electronic structure of films composed of various oxides of manganese. They showed changes in the structure of the L,,, edge of manganese, the K edge of oxygen and the shape of the low loss region for a series of Mn-0 films with compositions of MnO, Mn,O,, Mn,O, and MnO,. The lateral resolution of EELS is on the order of the probe size, which should be approximately 1 nm in a field emission gradient microscope for thin samples, e.g. in a 50 nm carbon film at 120 keV. In EDS the lateral resolution is poorer because of multiple scattering. As in EELS, it depends on the atomic number and thickness of the specimen. However, as shown by Kyser and Geiss [8] using Monte Carlo simulations, the most important factor in the determination of lateral resolution is the accelerating voltage. Monte Carlo techniques, now implemented on microcomputers, are being widely used to model the electron-matter scattering in thin films and, therefore, to predict the lateral resolution in EDS and EELS experiments. A suite of Monte Carlo programs developed by Joy [9] is available, free, through the software libraries of both EMSA and MAS. In an HRTEM study [lo] of alloy films of CoPt,, the influence of substrate temperature, during deposition, on the magnetic properties of the films was investigated. To correlate changes in the physical properties of the films with the magnetic data, the lattice parameter and grain size of the films were measured as a function of temperature over the range from room temperature (RT) to 600 ‘C. The lattice parameter was almost invariant over the temperature range studied. The grain size did show a slight increase in diameter from RT to 200 “C, and then had accelerated growth up to 600 “C. The scatter in the grain size was very small at low temperatures, but increased substantially above 200 “C following the pattern of increase in grain size. With HRTEM we observed the formation of some very small at temperatures above regions, 2-4 nm in diameter 300 “C (Fig. 2). Using the nano probe capabilities of the transmission electron microscope diffraction patterns were obtained from these regions, also shown in Fig. 2, which in many cases contained extra reflections which can be associated with an ordered phase in CoPt, This supports previous reports on the formation of an ordered phase of CoPt, above 200 “C [ 111. One of the classic problems in materials science which can be studied with HRTEM of thin films, SL films in plan view in this instance, is the measurement of elastic strain fields and lattice rotations associated with second phase or defect formation. Atomic
resolution images can be obtained from the structures associated with these strain fields. However, it is extremely difficult to make a reliable measurement of the positions of the atomic planes or columns. Dahmen et al. [ 121 developed a technique whereby a simulated ideal lattice structure with a slightly different lattice parameter is overlaid on the original micrographs. The resulting moire pattern reflects any differences in lattice spacings. Deviations from perfection in the original are seen as large changes in the moire pattern. Dahmen rt al. used this technique to visualize the strain field associated with 8’ precipitation in Al-Cu. Measurements of changes in the spacing and orientation of the moire pattern were used to quantify the distortion field at the end of the precipitate. In a somewhat similar application, Sinclair et al. [ 131 used optical microdiffraction to measure changes in lattice spacing associated with the spinodal decomposition of a Au-Ni alloy. The experiment consisted of using an optical bench to produce transforms from small areas of HRTEM micrographs of lattice images associated with the phase change. Local lattice parameters were then obtained by measurements off these transformed images. With the high resolution slow scan charge-coupled device (CCD) cameras now available for the transmission electron microscope, it is possible to perform this type of analysis without even taking a micrograph. With a CCD camera attached to a transmission electron microscope, a high resolution image can be acquired with a computer and the FFT analysis done on line at the microscope.
Fig. 2. HRTEM micrograph of a plan view of CoPt, alloy film deposited on a carbon substrate at 300 ‘C. The small, I-3 nm diameter, regions in the middle of large grains should be noted. These often had diffraction patterns (an example is shown) with extra reflections suggesting an ordered CoPt, structure. The large grains for the most part did not show superlattice reflections.
R. H. Geiss / Analytical electron microscopy
The imaging of magnetic domains in plan section SL films was first described by Hale et al. [ 141in 1959. The contrast arises as a result of the interaction of the incident electron beam with the magnetic induction in the sample. The interaction can be described in terms of the Lorentz force; thus the imaging of magnetic domains is known as Lorentz microscopy. An example of a Fresnel image is given in Fig. 3(a). Here the image is formed by defocusing the imaging lens. A Fresnel image delineates the domain walls in a magnetic thin film. The distribution of the electrons deflected by the magnetization in the sample can be imaged as a diffraction pattern in the back focal plane of the imaging lens. Using an aperture to exclude part of the pattern, which is analogous to dark field imaging in TEM, a blackand-white (BW) image, called a Foucault image, is
(4
(b) Fig. 3. (a) Phase contrast (Fresnel) Lorentz mlcrograph of magneuc domain walls in a Co-Pt alloy. (b) Diffraction contrast (Foucault) Lorentz micrograph from the same area showing the domains.
157
obtained. This image shows the BW magnetic domains characterizing the different magnetization directions in the sample (see Fig. 3(b)). McFadyen and Alexopoulos [ 151 quantified magnetic imaging using a scanning TEM technique called differential phase contrast (DPC). With DPC they can obtain vector maps detailing the magnetization directions in the sample. Cross-sectional views of SL films are usually not made unless one suspects there is an uneven elemental or defect distribution through the thickness of the film. However, considering the lateral resolution limitations of both EELS and EDS it might still be difficult to measure composition gradients in very thin films. Here, it might be necessary to make supplementary measurements with a depth profiling tool such as Rutherford backscattering. 3.2. Multilayer films Plan view TEM of ML films is usually very confusing since it is difficult to sort out the contribution from each of the layers, assuming they are sufficiently thin that a number of layers are included in a thinned plan view sample. An example of such a plan view is shown in Fig. 4, where Fig. 4(a) is a bright field micrograph of an ML sample with 6 layers. It is impossible to single out the contribution from any one layer in this micrograph. One feature frequently found in ML plan view imaging is a large number of moire fringes, as can be seen in this micrograph. One must be careful in deriving grain sizes from the moire images, however, as the precise interface where the fringes originated may not be easily established and there may be multiple overlapping of grains. If the crystallographic structure of the different layers is separable on a diffraction pattern, it may be possible to isolate a layer with dark field imaging. Also shown in Fig. 4(a) is the diffraction pattern from the ML structure. Prior knowledge of the phases of the layers allowed us to separate the pattern into b.c.c. and f.c.c. components. Using a small objective aperture it was possible to image the b.c.c. and f.c.c. components separately in dark field, as shown in Figs. 4(b) and 4(c) respectively. It is still impossible to identify clearly the individual layers and the buried interfaces. If the lattice parameters of the individual components are sufficiently different, it might be possible to separate phases. In this example, the b.c.c. pattern has a continuous diffuse ring suggesting that the grains are small and have a random orientation in the plane of the film. The small grain size can be seen in Fig. 4(b). On the contrary, the diffraction pattern from the f.c.c. structure shows at least two lattices with the same degree of preferred orientation which cannot be separated in dark field imaging. More commonly, ML films are studied in cross-section. Here the individual layers may be imaged and
information on each obtained if the resolution of the microscope permits. In Fig. 5 a high resolution image of a cross-section of the ML film used in Fig. 4 is given. It is easy to sort out the heavy metal bottom layer by the dark contrast of the 6 nm thick layer. However, it is impossible to delineate the next four layers, all f.c.c. and all containing transition elements, nickel, iron, copper and manganese, because the scattering from each of these atoms is essentially identical for electrons. The top layer, an oxidized heavy metal, is again easily seen. Careful inspection of Fig. 5 shows there are regions through the thickness of the transition metal layer zone where the lattice images are uniform throughout the thickness. In other regions there is a change in the imaging detail, but it is not sufficient to define the interfaces between the layers. The white dot contrast in some regions clearly defines some defect structures, but again does not provide any conclusive information on the interfaces.
Cc) Fig. 4. (a) Plan view of an ML film with six layers. Moire fringes are almost always present in ML films of this type since there will always be some regions with interfaces sufficiently aligned to produce them. The diffraction pattern shows both b.c.c. and f.c.c. patterns. The b.c.c. pattern is continuous and somewhat diffuse, while the f.c.c. patterns indicate some preferred orientation. The f.c.c. pattern shows that there are at least two different f.c.c. phases. (b) Dark field micrograph of (a) using the I IO b.c.c. diffraction ring. (c) Dark field micrograph of (a) using the 1I I and 200 f.c.c. reflections. The four arcs, two I I I and two 200, are too close to be individually selected with the objective aperture used for dark field imaging.
Fig. 5. HRTEM image of a cross-section through an ML alloy film like that in Fig. 4. The heavy metal layer, on the left, is easy to distinguish because of the greater scattering. The center of the film contains four f.c.c. layers, composed of similar transition metals, which are impossible to separate purely by inspection. The arrows at the top of the figure indicate the expected locations of the interfaces.
At present, we are evaluating the two techniques described earlier [ 12, 131, to determine the interfaces. In ML films where the components have widely differing atomic numbers such as in the Ti/W ML film studied by Dirks et al. [ 161, the contrast is is easily
R. H. Geiss / Analytical electron microscopy
The lateral resolution of EDS and EELS was described earlier. In most cases, for a cross-sectional sample 50 nm thick, the lateral resolution will be on the order of 2-4 nm. Therefore, neither EDS nor EELS will be sufficient to determine the structure of an interface if one wants the details on an atomic scale.
4. Summary Techniques of sample preparation continue to be refined. However, this aspect of the study of thin films is still often the most critical and demanding. Plan views are most likely to be used in the study of SL films, while cross-sectional views provide the most information in ML films. In both views all the analytical methods available in the microscope may be used in a complementary way for analysis, but in those cases where the details of the interface must be determined at the atomic level high resolution imaging provides the best tool. There is also a limit to the use of HRTEM; that is, if the atomic components of the two layers at the interface scatter electrons nearly identically, the contrast is too weak to allow the unique determination the location of the interface. To solve this problem, new methods of extracting information from high resolution images need to be developed.
Acknowledgments
The authors would like to thank R. F. C. Farrow, R. Marks, D. Bullock, D. K. Weller, J.-P. Nozieres, H. Lefakis, P. S. Alexopoulos and D. Dobbertin for
159
providing the samples used here and especially J. Baglin for his critical review of the manuscript.
References I J. C. Bravman and R. Sinclair, J. Electron Microsc. Tech., I (1984) 53. 2 J. P. Benedict, S. J. Klepis, W. G. Vandygrift and R. Anderson, Mater. Res. Sot. Symp. Proc., 199 (1990) 189. 3 S. J. Klepeis et al., Mater. Res. Sot. Symp. Proc., 115( 1988) 179. 4 R. H. Geiss and D. Bullock, unpublished. 5 R. H. Geiss and R. J. Savoy, in Howitt (ed.), Proc. Microbeam Analysis 1991, San Francisco Press, San Francisco, CA, 1991, p. 59. 6 R. H. Geiss, unpublished. 7 J. Patterson, 0. Krivanek and H. Poppa, Proc. EMSA 47th Anna. Meet., 1989. 8 D. F. Kyser and R. H. Geiss, in Ogilvie and Wittry (eds.), Proc. VIII Int. Co@ X-Ray Optics and Microanalysis, 1977, p. I IOa. 9 D. C. Joy, Monte Carlo simulations of electonsolid interactions, to be published. IO R. H. Geiss, R. F. C. Farrow, R. Marks and G. Harp, unpublished. II C. J. Chien, S. B. Clemens, S. B. Hagstrom, R. F. C. Farrow. C. H. Lee, E. E. Mariner0 and C. J. Lit-r. Mater. Res. Sot. Symp. Proc., 131 (1992) 465. I2 U. Dahmen, C. J. D. Hetherington and K. H. Westmacott, Proc. XIIth Ini. Cong. for Electron Microscopy, in Mater. Sci. 4 ( 1990) 338. I3 R. Sinclair, R. Gronsky and G. Thomas Acta Metall., 24 ( 1976) 789. I4 M. E. Hale, H. W. Fuller and H. Rubinstein, J. Appl. Phys., 30 ( 1959) 789. IS I. R. McFadyen and P. S. Alexopoulos, in Hadjipanyis and Prinz (eds.), Science and Technology of Nanostructural Magnetic Materia/s, Plenum New York, 1991, p. 99. 16 A. G. Dirks, A. M. Wolters and A. E. M. DeVersman, Thin Solid Films, 208 ( 1992) I8 I.
Analytical electron microscopy of thin films and interfaces R. H. Geiss
Abstract In this paper a brief overview is presented of some of the techniques available with transmission electron microscopy (TEM) to characterize thin films and interfaces. The critical role that sample preparation plays is discussed. Examples of the application of many of the analytical techniques used in TEM are given. These include nano area electron diffraction, energy-dispersive X-ray spectroscopy and electron energy loss spectroscopy. High resolution imaging is available at atomic resolution and new methods are being developed to analyze the images including computer simulation. Imaging of magnetic domains is important to the study of materials used in digital data storage.
1. Introduction
One of the most powerful techniques which can be used in the study of thin films and interfaces is high resolution transmission electron microscopy (HRTEM). When combined with analytical tools such as nano diffraction, energy loss spectroscopy (energydispersive spectroscopy (EDS)) and electron energy loss spectroscopy (EELS), it is given the name analytical electron microscopy (AEM). One of the major hurdles which must be overcome in any AEM investigation is the preparation of a thin sample which is representative of the material under investigation. It is appropriate that a brief discussion of the newer approaches to sample preparation be included. A few of the ways that AEM can be used to study thin films will be explored. Both single-layer (SL) and multilayer (ML) films in the plan and cross-section views will be considered. With modern AEM it is possible to obtain elemental data using EDS and EELS with a lateral resolution of 1- 10 nm depending on the microscope and the sample. Electron diffraction can be used to determine phase information, crystallographic orientation and lattice parameters from areas as small as l-2
nm in diameter
in microscopes
with
nano
probe
can be used to obtain images with better than 0.2 nm point-to-point resolution. To aid in the interpretation of these images computer simulation of the structure at the atomic scale should be included. Lorentz imaging of magnetic domains is very valuable in the study of digital storage media. To illustrate the above, examples will be given which show the use of AEM to determine (a) the grain size and details of the structure with high resolution capabilities.
HRTEM
imaging, (b) the structure of nano size regions in crystalline films with electron diffraction and (c) the use of EDS to determine the thickness of very thin films and also the minimum detection capabilities of EDS. Reference to the use of parallel EELS (PEELS) to determine the oxidation state of thin film oxides will be discussed. The use of Lorentz microscopy to image magnetic domain structures will be shown. Finally, some of the approaches used in the analysis of HRTEM lattice images of interface and defect structures will be discussed. The use of Fast Fourier Transforms (FFTs) and moire patterns may provide a mechanism for making measurements from these images.
2. Sample preparation The technique of specimen preparation of bulk materials has evolved from the methods of chemical and/or electrochemical thinning to techniques which include mechanical slicing, polishing and dimpling and a final thinning using argon ions. The latter methods have the advantage that they are more readily used on samples containing a combination of materials. It is almost impossible to find a chemical which would uniformly thin a sample consisting of an arbitrary mixture of components. The new techniques are also more easily used to prepare samples with localized defects or other important features. Mechanical polishing-ion thinning is the only method which can be used to prepare cross-sectional views reliably from samples on silicon or sapphire substrates. Although the development of these techniques has been particularly valuable in the study of semiconductor structures in cross-section, they can be used for a much broader variety of materials. The
Elsevier Sequoia
R. H. Geiss 1 Analytical
reader is referred to refs. 1 and 2 for a thorough discussion of these techniques. Two problems cause major concern, however; specimen heating which occurs during the many stages of the preparation, especially in the curing of the M Bond610, the epoxy recommended for silicon [ 11. The curing procedure for this epoxy suggests heating to 100 “C for 10 h or more. Heating may also occur during the ion thinning, especially if the thinning time is excessive, e.g. more than 12 h. There is also the possibility that during ion thinning some of the material which is removed may be redeposited onto other parts of the specimen. A technique in which samples are thinned by mechanical polishing only or with at most a few minutes of ion thinning [3] may help to alleviate the latter problem, but one must always be very aware of the potential problems associated with prolonged heating. For the preparation of plan section specimens any of the techniques discussed above can be used. However, if the effect of the substrate is not deemed critical, or if the heating is a problem, plan view samples of thin films can be prepared very quickly using a substrate of evaporated carbon on freshly cleaved mica. Very large area samples are obtained by releasing the film at the carbon-mica interface in water. If the exposure to water is undesirable, special thin window substrates such as silicon nitride on a silicon wafer can be used. Here the film is deposited on the Si,N, and used directly for TEM. The difficulty here is in manufacturing suitably thin silicon nitride windows. Films can be also prepared either by direct evaporation onto NaCl or by evaporation onto a substrate film previously deposited on NaCl. However, these films are also retrieved by floating them off in water. In principle, almost any sample can thinned for TEM study given the variety of techniques available today. However, one must be aware that sample preparation may be the most tedious and frustrating part of the study, especially in view of the criterion discussed above.
electron
microscop)
155
In a study of the structure of thin evaporated films of copper on carbon substrates [4] we found that films 1, 2 and 3 nm thick all had island-like structures consisting of small crystallites, 2-5 nm in dimension, surrounded by amorphous regions thought to be the carbon substrate (Fig. 1). However, an investigation of the 2 nm film with EDS using a 5 nm probe showed that the amorphous areas between the islands also contained copper. The thickness of the copper at the islands and at the amorphous areas between, as determined by the integrated intensity of the Cu Ka peak, varied from about 1 nm between the islands to more than 4 nm at the islands. The film thickness monitor and electron microprobe measurements had indicated that the average film thickness was right at the 2 nm mark. Even though one might argue the accuracy of the 1 and 4 nm thickness determinations, there is no doubt that the copper intensity on the islands was much greater than that from a large area sample and that there was copper between the islands. These measurements indicate a film profile which is quite different from that expected. Microdiffraction from areas between the islands showed that the structure was noncrystalline, suggesting that the regions between the islands may be amorphous. In an investigation using EDS to determine the minimum detectable limit (MDL) in thin films and alloys, Geiss and Savoy [5] showed that by using the total area under the peak, as the peak intensity, and the intensity of a single channel, as the background intensity, an MDL of about 150 ppm can be measured in most metal systems. A critical step in the analysis is the ability to extract a very small peak from a region of high background in the EDS spectra. This was done using a
3. Analysis 3.1. Single -layer films SL films may be single-element, compound, mixed element or alloy films. In the case of single-element films one is probably concerned with some aspect of the growth mechanism which might include the influence of the substrate, or with defects such as vacancies, dislocations or grain boundaries. In alloy SL films one can also study the structure and properties of different phases as well as the defects etc. To illustrate the use of AEM in the analysis of SL films consider the following examples.
Fig. 1. High resolution electron micrograph showing the island structure of a 2 nm thick copper film on a carbon substrate. EDS showed that the regions between the crystalline islands had as much as 1 nm of copper and that the islands were as thick as 4 nm.
personal computer and importing the EDS data in ASCII format into a spreadsheet template developed by one of these researchers [6]. Patterson et ul. [7] used PEELS to reveal differences in the electronic structure of films composed of various oxides of manganese. They showed changes in the structure of the L,,, edge of manganese, the K edge of oxygen and the shape of the low loss region for a series of Mn-0 films with compositions of MnO, Mn,O,, Mn,O, and MnO,. The lateral resolution of EELS is on the order of the probe size, which should be approximately 1 nm in a field emission gradient microscope for thin samples, e.g. in a 50 nm carbon film at 120 keV. In EDS the lateral resolution is poorer because of multiple scattering. As in EELS, it depends on the atomic number and thickness of the specimen. However, as shown by Kyser and Geiss [8] using Monte Carlo simulations, the most important factor in the determination of lateral resolution is the accelerating voltage. Monte Carlo techniques, now implemented on microcomputers, are being widely used to model the electron-matter scattering in thin films and, therefore, to predict the lateral resolution in EDS and EELS experiments. A suite of Monte Carlo programs developed by Joy [9] is available, free, through the software libraries of both EMSA and MAS. In an HRTEM study [lo] of alloy films of CoPt,, the influence of substrate temperature, during deposition, on the magnetic properties of the films was investigated. To correlate changes in the physical properties of the films with the magnetic data, the lattice parameter and grain size of the films were measured as a function of temperature over the range from room temperature (RT) to 600 ‘C. The lattice parameter was almost invariant over the temperature range studied. The grain size did show a slight increase in diameter from RT to 200 “C, and then had accelerated growth up to 600 “C. The scatter in the grain size was very small at low temperatures, but increased substantially above 200 “C following the pattern of increase in grain size. With HRTEM we observed the formation of some very small at temperatures above regions, 2-4 nm in diameter 300 “C (Fig. 2). Using the nano probe capabilities of the transmission electron microscope diffraction patterns were obtained from these regions, also shown in Fig. 2, which in many cases contained extra reflections which can be associated with an ordered phase in CoPt, This supports previous reports on the formation of an ordered phase of CoPt, above 200 “C [ 111. One of the classic problems in materials science which can be studied with HRTEM of thin films, SL films in plan view in this instance, is the measurement of elastic strain fields and lattice rotations associated with second phase or defect formation. Atomic
resolution images can be obtained from the structures associated with these strain fields. However, it is extremely difficult to make a reliable measurement of the positions of the atomic planes or columns. Dahmen et al. [ 121 developed a technique whereby a simulated ideal lattice structure with a slightly different lattice parameter is overlaid on the original micrographs. The resulting moire pattern reflects any differences in lattice spacings. Deviations from perfection in the original are seen as large changes in the moire pattern. Dahmen rt al. used this technique to visualize the strain field associated with 8’ precipitation in Al-Cu. Measurements of changes in the spacing and orientation of the moire pattern were used to quantify the distortion field at the end of the precipitate. In a somewhat similar application, Sinclair et al. [ 131 used optical microdiffraction to measure changes in lattice spacing associated with the spinodal decomposition of a Au-Ni alloy. The experiment consisted of using an optical bench to produce transforms from small areas of HRTEM micrographs of lattice images associated with the phase change. Local lattice parameters were then obtained by measurements off these transformed images. With the high resolution slow scan charge-coupled device (CCD) cameras now available for the transmission electron microscope, it is possible to perform this type of analysis without even taking a micrograph. With a CCD camera attached to a transmission electron microscope, a high resolution image can be acquired with a computer and the FFT analysis done on line at the microscope.
Fig. 2. HRTEM micrograph of a plan view of CoPt, alloy film deposited on a carbon substrate at 300 ‘C. The small, I-3 nm diameter, regions in the middle of large grains should be noted. These often had diffraction patterns (an example is shown) with extra reflections suggesting an ordered CoPt, structure. The large grains for the most part did not show superlattice reflections.
R. H. Geiss / Analytical electron microscopy
The imaging of magnetic domains in plan section SL films was first described by Hale et al. [ 141in 1959. The contrast arises as a result of the interaction of the incident electron beam with the magnetic induction in the sample. The interaction can be described in terms of the Lorentz force; thus the imaging of magnetic domains is known as Lorentz microscopy. An example of a Fresnel image is given in Fig. 3(a). Here the image is formed by defocusing the imaging lens. A Fresnel image delineates the domain walls in a magnetic thin film. The distribution of the electrons deflected by the magnetization in the sample can be imaged as a diffraction pattern in the back focal plane of the imaging lens. Using an aperture to exclude part of the pattern, which is analogous to dark field imaging in TEM, a blackand-white (BW) image, called a Foucault image, is
(4
(b) Fig. 3. (a) Phase contrast (Fresnel) Lorentz mlcrograph of magneuc domain walls in a Co-Pt alloy. (b) Diffraction contrast (Foucault) Lorentz micrograph from the same area showing the domains.
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obtained. This image shows the BW magnetic domains characterizing the different magnetization directions in the sample (see Fig. 3(b)). McFadyen and Alexopoulos [ 151 quantified magnetic imaging using a scanning TEM technique called differential phase contrast (DPC). With DPC they can obtain vector maps detailing the magnetization directions in the sample. Cross-sectional views of SL films are usually not made unless one suspects there is an uneven elemental or defect distribution through the thickness of the film. However, considering the lateral resolution limitations of both EELS and EDS it might still be difficult to measure composition gradients in very thin films. Here, it might be necessary to make supplementary measurements with a depth profiling tool such as Rutherford backscattering. 3.2. Multilayer films Plan view TEM of ML films is usually very confusing since it is difficult to sort out the contribution from each of the layers, assuming they are sufficiently thin that a number of layers are included in a thinned plan view sample. An example of such a plan view is shown in Fig. 4, where Fig. 4(a) is a bright field micrograph of an ML sample with 6 layers. It is impossible to single out the contribution from any one layer in this micrograph. One feature frequently found in ML plan view imaging is a large number of moire fringes, as can be seen in this micrograph. One must be careful in deriving grain sizes from the moire images, however, as the precise interface where the fringes originated may not be easily established and there may be multiple overlapping of grains. If the crystallographic structure of the different layers is separable on a diffraction pattern, it may be possible to isolate a layer with dark field imaging. Also shown in Fig. 4(a) is the diffraction pattern from the ML structure. Prior knowledge of the phases of the layers allowed us to separate the pattern into b.c.c. and f.c.c. components. Using a small objective aperture it was possible to image the b.c.c. and f.c.c. components separately in dark field, as shown in Figs. 4(b) and 4(c) respectively. It is still impossible to identify clearly the individual layers and the buried interfaces. If the lattice parameters of the individual components are sufficiently different, it might be possible to separate phases. In this example, the b.c.c. pattern has a continuous diffuse ring suggesting that the grains are small and have a random orientation in the plane of the film. The small grain size can be seen in Fig. 4(b). On the contrary, the diffraction pattern from the f.c.c. structure shows at least two lattices with the same degree of preferred orientation which cannot be separated in dark field imaging. More commonly, ML films are studied in cross-section. Here the individual layers may be imaged and
information on each obtained if the resolution of the microscope permits. In Fig. 5 a high resolution image of a cross-section of the ML film used in Fig. 4 is given. It is easy to sort out the heavy metal bottom layer by the dark contrast of the 6 nm thick layer. However, it is impossible to delineate the next four layers, all f.c.c. and all containing transition elements, nickel, iron, copper and manganese, because the scattering from each of these atoms is essentially identical for electrons. The top layer, an oxidized heavy metal, is again easily seen. Careful inspection of Fig. 5 shows there are regions through the thickness of the transition metal layer zone where the lattice images are uniform throughout the thickness. In other regions there is a change in the imaging detail, but it is not sufficient to define the interfaces between the layers. The white dot contrast in some regions clearly defines some defect structures, but again does not provide any conclusive information on the interfaces.
Cc) Fig. 4. (a) Plan view of an ML film with six layers. Moire fringes are almost always present in ML films of this type since there will always be some regions with interfaces sufficiently aligned to produce them. The diffraction pattern shows both b.c.c. and f.c.c. patterns. The b.c.c. pattern is continuous and somewhat diffuse, while the f.c.c. patterns indicate some preferred orientation. The f.c.c. pattern shows that there are at least two different f.c.c. phases. (b) Dark field micrograph of (a) using the I IO b.c.c. diffraction ring. (c) Dark field micrograph of (a) using the 1I I and 200 f.c.c. reflections. The four arcs, two I I I and two 200, are too close to be individually selected with the objective aperture used for dark field imaging.
Fig. 5. HRTEM image of a cross-section through an ML alloy film like that in Fig. 4. The heavy metal layer, on the left, is easy to distinguish because of the greater scattering. The center of the film contains four f.c.c. layers, composed of similar transition metals, which are impossible to separate purely by inspection. The arrows at the top of the figure indicate the expected locations of the interfaces.
At present, we are evaluating the two techniques described earlier [ 12, 131, to determine the interfaces. In ML films where the components have widely differing atomic numbers such as in the Ti/W ML film studied by Dirks et al. [ 161, the contrast is is easily
R. H. Geiss / Analytical electron microscopy
The lateral resolution of EDS and EELS was described earlier. In most cases, for a cross-sectional sample 50 nm thick, the lateral resolution will be on the order of 2-4 nm. Therefore, neither EDS nor EELS will be sufficient to determine the structure of an interface if one wants the details on an atomic scale.
4. Summary Techniques of sample preparation continue to be refined. However, this aspect of the study of thin films is still often the most critical and demanding. Plan views are most likely to be used in the study of SL films, while cross-sectional views provide the most information in ML films. In both views all the analytical methods available in the microscope may be used in a complementary way for analysis, but in those cases where the details of the interface must be determined at the atomic level high resolution imaging provides the best tool. There is also a limit to the use of HRTEM; that is, if the atomic components of the two layers at the interface scatter electrons nearly identically, the contrast is too weak to allow the unique determination the location of the interface. To solve this problem, new methods of extracting information from high resolution images need to be developed.
Acknowledgments
The authors would like to thank R. F. C. Farrow, R. Marks, D. Bullock, D. K. Weller, J.-P. Nozieres, H. Lefakis, P. S. Alexopoulos and D. Dobbertin for
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providing the samples used here and especially J. Baglin for his critical review of the manuscript.
References I J. C. Bravman and R. Sinclair, J. Electron Microsc. Tech., I (1984) 53. 2 J. P. Benedict, S. J. Klepis, W. G. Vandygrift and R. Anderson, Mater. Res. Sot. Symp. Proc., 199 (1990) 189. 3 S. J. Klepeis et al., Mater. Res. Sot. Symp. Proc., 115( 1988) 179. 4 R. H. Geiss and D. Bullock, unpublished. 5 R. H. Geiss and R. J. Savoy, in Howitt (ed.), Proc. Microbeam Analysis 1991, San Francisco Press, San Francisco, CA, 1991, p. 59. 6 R. H. Geiss, unpublished. 7 J. Patterson, 0. Krivanek and H. Poppa, Proc. EMSA 47th Anna. Meet., 1989. 8 D. F. Kyser and R. H. Geiss, in Ogilvie and Wittry (eds.), Proc. VIII Int. Co@ X-Ray Optics and Microanalysis, 1977, p. I IOa. 9 D. C. Joy, Monte Carlo simulations of electonsolid interactions, to be published. IO R. H. Geiss, R. F. C. Farrow, R. Marks and G. Harp, unpublished. II C. J. Chien, S. B. Clemens, S. B. Hagstrom, R. F. C. Farrow. C. H. Lee, E. E. Mariner0 and C. J. Lit-r. Mater. Res. Sot. Symp. Proc., 131 (1992) 465. I2 U. Dahmen, C. J. D. Hetherington and K. H. Westmacott, Proc. XIIth Ini. Cong. for Electron Microscopy, in Mater. Sci. 4 ( 1990) 338. I3 R. Sinclair, R. Gronsky and G. Thomas Acta Metall., 24 ( 1976) 789. I4 M. E. Hale, H. W. Fuller and H. Rubinstein, J. Appl. Phys., 30 ( 1959) 789. IS I. R. McFadyen and P. S. Alexopoulos, in Hadjipanyis and Prinz (eds.), Science and Technology of Nanostructural Magnetic Materia/s, Plenum New York, 1991, p. 99. 16 A. G. Dirks, A. M. Wolters and A. E. M. DeVersman, Thin Solid Films, 208 ( 1992) I8 I.