microscopia de transmissão

microscopia de transmissão

(Parte 1 de 3)

nature materials | VOL 8 | APRIL 2009 | w.nature.com/naturematerials 263 insight | review articles Published online: 24 march 2009 | doi: 10.1038/nmat2380

Problem solving at the atomic scale has become an essential part of both fundamental science and modern device development. Hand in hand with the desire and ability to control materials at the level of atoms, comes a need to image and quantify materials properties with atomic resolution1. The success or failure of semiconductor devices and their continued miniaturization relies on critical components. These include the 5–10-atom-thick gate oxides in transistors2–4, magnetoresistive read heads in hard drives that have layers which are only 1–2 nm thick5–7 and tunnel junctions in magnetic memories that are also of comparable thickness8–10. The challenge in all these systems is not only to determine the thickness and composition of the ‘bulk’ portion of these ultrathin layers, but also the chemistry, interdiffusion and electronic structure of the ‘buried’ interfaces between the layers.

Although no single method can yet obtain complete information about buried structures at the atomic scale, this goal is being approached from several directions. This includes compositional (but not bonding) information from atom-probe microscopy, structural information from transmission electron microscopy and electron tomography (see the Review by Midgley and Dunin-Borkowski elsewhere in this Insight11) and scanning transmission electron microscopy (STEM), which will be discussed here. STEM has proved very effective in measuring not only compositional changes at buried interfaces, but also the electronic structure and bonding that have relevance for the mechanical12–14 and transport properties2,15 of a device. Detection of single dopant atoms with STEM, both on free surfaces16 and buried inside devices17,18, has proved useful in studies ranging from the characterization of catalysts19,20 to understanding the materials limits for transistor scaling17. It is also possible to detect and image the spatial distribution of single vacancies — either directly21,2, or by their strain fields, or spectroscopically from their electronic fingerprints on the local densities of states23.

The basic concept of a STEM is relatively simple — a high-energy electron beam is focused down to a small spot and fired through a thin sample. Signals from scattered electrons and ionized atoms are recorded as the beam is scanned across the sample to build up a twodimensional map. Chemical and bonding information along each projected atomic column can be obtained by measuring the energy lost by transmitted electrons to core and valence excitations at each point on the map (Fig. 1). The performance of the instrument is determined by how small a spot the electron beam can be focused to, and how much current can be maintained in that beam. Recent advances in the control of lens aberrations through the successful structure and bonding at the atomic scale by scanning transmission electron microscopy david a. muller1

A new generation of electron microscopes is able to explore the microscopic properties of materials and devices as diverse as transistors, turbine blades and interfacial superconductors. All of these systems are made up of dissimilar materials that, where they join at the atomic scale, display very different behaviour from what might be expected of the bulk materials. Advances in electron optics have enabled the imaging and spectroscopy of these buried interface states and other nanostructures with atomic resolution. Here I review the capabilities, prospects and ultimate limits for the measurement of physical and electronic properties of nanoscale structures with these new microscopes.

development of multipole-based aberration correctors24–27 have enabled sub-ångström beam sizes28–30 and rapid, atomic-resolution compositional imaging31,32. Figure 2 shows that progress in improving spatial resolution has been both rapid and ongoing, so that 0.5-Å beam sizes are now possible. (Note that for a limit of resolution δx, the resolving power of the microscope33 is defined as the reciprocal distance, 1/δx).

Given that interatomic distances are all larger than 0.5 Å, one might well suspect, as did Dennis Gabor in 1948, that “resolution will have to stop here for a lack of objects” (ref. 34). Although this lack of suitable test objects does make it difficult to test the resolution of a microscope, there is no shortage of problems that could be tackled with smaller beam sizes. Atoms in amorphous materials or tilted crystals can appear much closer when viewed in projection. Being able to resolve arbitrarily short distances at multiple projections could lead to direct structural solutions for general amorphous materials, or atomic-resolution reconstructions of three-dimensional objects by electron tomography. The number of School of Applied and Engineering Physics, Cornell University, Ithaca, New York 14853, USA. e-mail: dm24@cornell.edu ab cd

Figure 1 | Compositional imaging at the atomic scale. a–d, An atomic- resolution chemical map of a (La,Sr)MnO/SrTiO multilayer showing the La–M (a), Ti–L (b) and Mn–L (c) edge spectroscopic signals merged into a composite map (d). The white circles show the location of the columns of La atoms. The spectral series was recorded on an aberration-corrected scanning electron microscope using electron energy-loss spectroscopy. The data took less than 1 min to acquire. (From ref. 32; © 2008 AAAS.)

264 nature materials | VOL 8 | APRIL 2009 | w.nature.com/naturematerials review articles | insightNAture mAterIAls doi: 10.1038/nmat2380 resolvable projections increases roughly as the square of the resolution, so even improvements in resolution by a factor of two or three could make a dramatic difference. Even when such small distances do not need to be measured, a smaller probe size allows for higher contrast images and improved detectability of single dopant atoms, enabling studies of a wider range of dopant species and thicker, more realistic, substrates.

This article will explore the limits and opportunities for the new generation of aberration-corrected microscopes. After a brief description of the STEM itself, the improvements in resolution and current due to aberration-corrected optics and improved electron sources are detailed. Improved energy resolution from monochromators and higher-resolution spectrometers has enabled nanoscale measurements of the photonic properties of materials and waveguides. Finally, with the improvements in instrumentation, it is the sample itself and the radiation dose it can tolerate that seems now to be the limit to signal and resolution.

the scanning transmission electron microscope The first STEM to include the necessary field-emission source and high-vacuum needed to keep samples clean during irradiation was developed by Crewe in 1964 (ref. 35), and within a few years they were demonstrating images of single heavy atoms15 and their diffusion across thin carbon films36. The popular dark-field and highangle dark-field imaging modes were demonstrated early during the machine’s development, including in the collection of lattice images from nanoparticles37. As the original microscopes were intended for biological research, these samples were used more as resolution tests than for materials studies — with one exception being the study of catalyst particles19. Adoption of the STEM technique was slow owing to difficulties in sample preparation of bulk materials and a lack of widespread instrument availability as a result of the vacuum and electronics requirements.

Almost a decade later, with the development of a commercially manufactured STEM and better specimen preparation tools such as ion-milling, there was a renewed interest in the instrument as an analytical tool for metallurgy and in the semiconductor industry. Some of these second-generation instruments are still in operation today. In the mid-1990s, a new generation of commercial fieldemission TEMs with scanning attachments made STEM techniques widely available to the materials research community and industry.

The major components of a STEM are shown in Fig. 3. A field-emission gun provides a high-coherence source of electrons that is accelerated to between 100 and 300 keV — energies sufficient to penetrate samples up to 100 nm thick without significant beam spreading. A series of electron lenses and corrective optics placed before the sample focuses the beam down to a diameter smaller

Light microscope

Electron microscope

Aberration-corrected electron microscope

Zach, Haider (1 kV SEM)


Amici Abbe

Ruska (75 kV)

Borries and Ruska Marton, Ardenne (100 kV)

Dietrich (200 kV)



Haider (200 kV)

F eatur e siz

Resolving po

Figure 2 | Hardware advances in imaging microscopies. Optical microscopy reached its far-field diffraction limit roughly a century ago. Electron microscopy, exploiting the reduction in electron wavelength with increasing beam voltage, showed steady increases in resolving power for over 50 years, until limited by radiation damage. Recent developments in correcting electron-optical aberrations have led to improvements in resolving power without having to increase the electron beam energy. Although still far from their ultimate diffraction limits, corrected electron microscopes have already demonstrated the ability to resolve subångström image features. (Adapted from ref. 114; © 2009 Elsevier.)

Source Lenses


ADF detector

Electron spectrometer

Figure 3 | major elements of a scanning transmission electron microscope. A high-brightness electron source produces a 100–300- keV electron beam with an energy spread of 0.3–1 eV, which can be narrowed to below 100 meV with a monochromator. Round magnetic lenses and corrective multipole optics focus the beam to a spot size of between 0.05 and ~0.3 nm, which is scanned across an electron-transparent sample. To a first approximation, when the beam is placed on an atom column, strong Rutherford-like scattering deflects the transmitted electrons to form brighter features in an annular dark-field (ADF) image, with less scattering between the columns. Inelastic scattering is strongly peaked in the forward direction and is collected simultaneously with the ADF signal. The energy losses of the transmitted electrons reflect characteristic excitations of the sample in a frequency range spanning the near-infrared to hard X-ray regions, allowing electronic and elemental identification from a single column of atoms.

nature materials | VOL 8 | APRIL 2009 | w.nature.com/naturematerials 265 insight | review articlesNAture mAterIAls doi: 10.1038/nmat2380 than the spacings between the projected atomic columns. The beam is then scanned across the sample, and the differences in electronscattering when the beam is located on and off the atomic columns are recorded. The optimal sample thickness is a balance between being thin enough that the electron probe will not spread significantly in the sample, and thick enough so that bulk-like, and not surface, atoms dominate the signal.

Elastically scattered electrons can undergo large angle deflections, and hence an annular dark-field (ADF) detector is placed below the sample to collect the scattered electrons. The high-angle electronscattering cross-section scales roughly as the atomic number Z1.7, asymptotically tending towards the Rutherford scattering limit of Z2, and so this imaging mode is sometimes called ‘Z-contrast’ imaging18,37–39. Provided the collection angles are three or more times larger than the probe convergence angle, the incoherent imaging approximation is satisfied for thin samples, thereby allowing relatively straightforward image interpretation39–42. To avoid contrast reversals with sample thickness, the detector also needs to be sufficiently large to avoid any strongly diffracted beams from being collected19,38,43. When this condition is just met, diffuse scattering from phonons and lattice distortions can provide significant image contrast. This low-angle ADF signal is useful for strain mapping as well as for imaging light dopant atoms and vacancies23,43,4. At even higher angles, these effects become less pronounced and the atomicnumber contrast from nuclear scattering dominates. The resulting high-angle ADF images are easier to interpret, with the brighter features resulting from high-Z atoms, or higher densities of atoms. Electron channelling in the crystal can still provide unexpected contrast that may confound a direct quantitative analysis, but generally the contrast is still sufficient that model inputs for quantitative simulations can be directly determined39,43,45–47.

Compositional and bonding information can be obtained from electron energy-loss spectroscopy (EELS) analysis of the inelastically scattered electrons. The time-varying, electric-field pulses from the incident-beam electrons transfer energy to the sample over a range of frequencies, from the infrared to the X-ray regime, as they pass near target atoms. The resulting core-level excitations provide unique spectroscopic information about the excited atom and its bonding states (Fig. 4). Fortunately, the inelastic scattering is strongly peaked in the forward direction and so passes through the hole in the centre of the ADF detector. A spectrometer placed on axis can collect upwards of 90% of the EEL signal without compromising the ADF signal. This ADF/EELS geometry has been used since the earliest days of STEM, where the larger ADF signal is used to form an image and locate the probe for selected EELS measurements35,37. By 1986, 0.4-nm resolution composition profiles were demonstrated by Scheinfein and Isaacson48. Development of low-noise parallel detectors enabled the recording of spectra that showed details of chemical bonding from selected points in an atomic-resolution image49.

The time required to record a single spectrum using that technology was ~1–10 s. Even a spectral image with as few as 64 × 64 pixels could take hours to record — a serious challenge when typical sample drift rates are ~5 Å min−1. Despite this, atomic-resolution composition maps have been demonstrated (after significant drift correction and data processing)50. One of the key benefits of aberration correctors is that they increase the signal, allowing collection times to be reduced from hours to under a minute (Fig. 1), and also increase the signal-to-noise ratio so that bonding features are visible32. A full twodimensional chemical map obtained by EELS that also contains bonding information requires yet another order of magnitude increase in the beam current beyond that needed for simple compositional imaging in order to improve the signal-to-noise ratio. This has been done by further correcting the electron optical aberrations to fifthorder, rather than only to third-order as in earlier correctors. With the aberrations corrected over a wider angular range, the numerical aperture of the probe-forming lens can then be increased, allowing more electrons to be brought to a focused spot. It has also proved necessary to correct the collection optics so that the additional electrons input at the larger angles could still be collected31,32.

A very powerful feature of EELS is that it provides both compositional and bonding information at very high spatial resolution. Figure 4 shows the EELS excitation process whereby the incident electron transfers energy to the target atom, exciting a core electron to empty states above the Fermi level (EFermi). The energy loss of the transmitted electron is recorded by the spectrometer and provides information equivalent to X-ray absorption spectroscopy, where the electron’s momentum transfer has the same role as the polarization of the X-ray beam51,52. The core-level binding energy that marks the EELS edge onset allows a unique elemental identification to be made. The shape of the edge itself reflects the underlying local partial density of states modified by the presence of a core hole53–57. The core hole is relatively well screened in metals, allowing the spectrum to be interpreted in terms of the ground state local density of


Dipole transition

Core level

Figure 4 | electron energy-loss spectroscopy (eels) as a probe of local electronic structure and composition. EELS records the energy lost by an incident electron as it excites a core electron to unoccupied local states, leaving behind a core hole. Analysing the resulting loss spectrum of the transmitted beam provides information on the unoccupied local density of states, partitioned by site, chemical species and angular momentum. The site specificity is achieved by forming a probe small enough to only excite electrons on a single atomic column. The chemical specificity is obtained from the uniqueness of the core-level binding energy. The angular momentum of the final state is determined by the dipole selection rules and the angular momentum of the initial state. When the core hole is well screened, the shape of the EELS spectrum reflects the ground state density of states. In systems with strong core-hole coupling, excitonic effects can dramatically alter the shape of spectrum, but also provide useful local fingerprints of the formal charge and coordination chemistry.

266 nature materials | VOL 8 | APRIL 2009 | w.nature.com/naturematerials review articles | insightNAture mAterIAls doi: 10.1038/nmat2380 states51,54–56. Core-hole effects can dominate in insulators, especially on the cation sites, contracting the final state wavefunctions and allowing a more molecular or crystal-field-like interpretation of the local bonding environment to be made51,56,58–60.

The microscopic insight into electronic structure at buried interfaces provided by EELS has proved useful in the study of transistor miniaturization. The narrowest feature in modern transistors is the gate oxide. This is a dielectric layer roughly 5–6 atoms thick and tra- ditionally made from SiO2 — which was just recently replaced with Hf-based dielectric materials that for the moment are about twice as thick3,61. As device-scaling continues, these layers will also have to shrink. However, their continued scaling is limited by the interfacial reaction layers and altered electronic structure present at the interface2,62. For example, in the five-Si-atom-thick gate oxide shown in Fig. 5, at least two of those five atoms will be at the silicon/oxide interfaces, and EELS shows a very different local electronic structure for the interfacial atoms. These atoms are expected to have very different electrical and optical properties from the desired bulk SiO2, yet they comprise a significant fraction of the dielectric layer2,57.

In more-ionic materials, the local EELS fingerprint can be used to map the spatial distribution of formal charges at interfaces and defects63–65. The 3d transition metal edges are well suited to this task and atomic-scale analysis of the Ti–L edge has provided the spatial distribution of conduction electrons and their screening lengths in

LaTiO3/SrTiO3 multilayers, where a highly conducting sheet of electrons is formed at the interface between a Mott insulator and a band insulator64. Corrected microscopes should prove especially useful for identifying single dopant atoms inside a specific device, where the ‘tyranny of small numbers’ requires two-dimensional imaging instead of simply line profiles to obtain sufficient counting statistics from volumes that contain only a few atoms47,6.

advances in electron optics Although the ability to focus an electron beam down to a spot size of around 2 Å with a simple, round magnetic lens is remarkable, 2 Å is still far from the fundamental performance limit that could be possible for electron imaging. The typical electron wavelength in a TEM is on the order of 0.02 Å, yet the spatial resolution of such a microscope is 1–2 Å, between 50 and 100 times ‘worse’ than the wavelength. Wavelength for wavelength, conventional electron lenses have similar optical properties to the bottom of a beer bottle with respect to spherical and chromatic aberrations. This is not the result of imperfect engineering, but rather a consequence of Laplace’s equation, where the electric and magnetic potentials used to focus the electron beam cannot take on arbitrary shapes in free space. The result is that positive spherical and chromatic aberrations are unavoidable for static round lenses. This fundamental limit was first identified by Scherzer67 in 1936, who also suggested methods to correct these aberrations by considering non-static, non-round lenses or placing free charges in the path of the beam68. Of these, the most practical approach has proved to be abandoning cylindrical symmetry for some of the lenses and instead adding a series of multipole optical elements that correct the leading terms in an everincreasing power-series expansion of the wavefront distortions69.

It took nearly 50 years before the precision, stability and software control needed to implement these corrector designs could be realized24,25,70. Figure 2 shows that since the proof of concept was established, corrector development has been rapid (and is still continuing), with probe sizes close to 0.5 Å now reached29,30. Aberration correctors have now removed spherical aberration (and higher-order geometric aberrations) as limiting factors24,26–28,31. Instead, chromatic aberration, combined with the finite energy spread of the electron source, are the current limits on spatial resolution27,31. The resolution and aperture size can be increased by building a chromatic aberration corrector71, increasing the beam voltage or reducing energy spread of the source. Improvements of a factor of two appear possible, but not without cost: increasing the beam voltage also increases the risk of radiation damage — so beam voltage is more likely to be decreased in future instruments, only exacerbating the chromatic problems. Reducing the source energy spread by adding a monochromator reduces the available beam current. Cryogenically cooled sources can narrow the source energy spread and improve gun brightness72, but have proved difficult to implement owing to the required mechanical stability.

Corrective optics are also needed to retain and improve the collection efficiency of the EELS spectrometer31,32. Without these corrections, much of the additional current provided by the probe corrector will not enter the spectrometer, yet will add to the radiation dose of the sample. The effects can be dramatic — an uncorrected spectrometer on a corrected microscope can lose as much as 90% of the inelastic signal and provide no more signal than a completely uncorrected microscope. These very low collection efficiencies introduce additional problems in image formation and interpretation, because the incoherent imaging approximation no longer holds73–75. Diffraction contrast artefacts, such as bend contours and thickness fringes, can appear and elastic scattering artefacts can dominate the EELS spectrum image76–78. Strong elastic scattering on heavy-atom columns can prevent inelastic electrons from entering the spectrometer, resulting in contrast reversals and inelastic images that reflect mostly the elastic scattering and not the elemental distribution in the sample.

A common symptom of this elastic scattering artefact is that all inelastic maps peak at the same spatial positions, independent of the actual location of the elements. This problem was identified and demonstrated by quantitative theory and experimental work for Bi0.5Sr0.5MnO3 — where the ‘inelastic’ images are dominated by the strong elastic scattering from the Bi sites that largely prevents electrons on the Bi column from entering the detector76. The result for the 〈1〉 zone axis is an apparent map of Mn, where the Mn signal is smallest on the only columns that actually contain Mn, and largest where there are no Mn atoms76. The effect can be suppressed by increasing the collection angle so that almost all of the elastically

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