scanning transmission electron microscopy
Advances in scanning
transmission electron
microscopy
Stephen Pennycook, Scanning Transmission Electron Microscopy Group, Materials Science
and Technology Division, Oak Ridge National Laboratory, TN, USA
The Pioneer Years
The scanning transmission electron microscope
(STEM) was invented by Baron Manfred von
Ardenne [1, 2] not long after Max Knoll and
Ernst Ruska invented the first transmission
electron microscope (TEM) [3], and immediately
highlighted the key problem with the technique,
that of noise. For high resolution, small probes are
needed and the available current is limited. Von
Ardenne rapidly abandoned the STEM in favor
of Ruska’s TEM design which illuminates a large
area, avoiding the issue. It was only when Albert
Crewe incorporated a cold field-emission source
into the instrument that it became a viable form
of microscope [4, 5]. His field-emission source
had orders of magnitude higher brightness, and
probes approaching atomic dimensions could
carry enough current to form acceptable images.
Furthermore, the resolution-controlling optics
are before the specimen in the STEM, avoiding
many of the chromatic aberration issues in the
objective lens, and Crewe also took advantage of
the flexible detection optics. Crewe’s STEM was
a huge breakthrough; it was the first electron
microscope to resolve single atoms [6, 7] (the fieldion microscope imaged atoms earlier [8]). Crewe’s
Chicago group dominated the field, generating
a large body of very impressive research results,
not only the first single atom imaging, but the
introduction of the annular detector, the discrimination of atoms of different atomic number (Z)
[9], and the first simultaneous annular dark field
(ADF) imaging and electron energy loss spectroscopy (EELS) [10], among many others. For a
comprehensive review of the activity of the Crewe
laboratory see the account by Peter Hawkes [11].
These spectacular breakthroughs stimulated the
introduction of a commercial dedicated STEM
with a cold field-emission gun by VG Microscopes, which allowed many other laboratories
to experiment with the instrument, notably
Mick Brown and Archie Howie at the Cavendish
Laboratory, University of Cambridge, UK [12,
13], Christian Colliex at the Université ParisSud, France [14, 15] and John Cowley at Arizona
State University, USA [16, 17]. The author was
fortunate enough to be a graduate student,
then a postdoctoral scientist at the Cavendish
Laboratory with the second VG Microscopes HB5
ever built (the first going to Ron Burge at the
University of London). While pioneering for the
time, with their UHV cold field-emission gun
and high vacuum, bakeable specimen chamber,
digital electronics was in its infancy. Images were
recorded on Polaroid or 35-mm film, and EELS
data were typically recorded on a chart recorder.
Figure 1 shows one of the first EELS spectra to
be recorded digitally using a modified Link
Systems X-ray multichannel analyzer. The digital
acquisition allowed spectra on and off the defect
to be subtracted, the first example of a spatial
difference technique in EELS [18]. It provided
an accurate background subtraction for the N
K edge despite the presence of strong extended
fine structure variations on the tail of the carbon
K-loss signal [19].
Z-contrast for materials science
While the Crewe group was primarily interested
in biological applications, the other groups had
interests in materials science, samples that were
typically crystalline. The Z-contrast technique
originally introduced by Crewe and coworkers
used a ratio between the ADF image (dominated by elastic scattering) and the EELS image
(inelastic scattering), which for single atoms
had a ratio of approximately Z [20]. However, in
crystalline materials diffraction contrast usually
dominated such Z-contrast, and the ratio method
did not work reliably. The answer was to increase
the inner angle of the ADF detector to avoid the
domination by coherent Bragg diffraction [12,
21]; the high-angle annular dark field (HAADF)
technique was born.
On moving to Oak Ridge National Laboratory
(ORNL) in 1982 I applied the HAADF technique
to ion-implanted Si, showing how images of
dopant distributions could be obtained avoiding
strong diffraction contrast [22], and pondered
what such an image would look like if the STEM
were capable of atomic resolution, would it show
incoherent characteristics? Cowley had already
obtained a lattice image using a wide-angle,
Crewe-style ADF detector [23], although he did
not mention any incoherent characteristics, for
example, the freedom from contrast reversals
which were implicit in the Crewe single atom images, and which had been predicted theoretically
for single atoms by Andreas Engel [24]. The question was answered in 1988 when VG Microscopes
delivered an HB501UX to ORNL equipped with
a high-resolution pole piece. Images of the high-
temperature superconductors YBa2Cu3O7-x and
ErBa2Cu3O7-x showed both the expected Z-contrast
and no contrast reversal with objective lens focus
[25]. Images of Si and InP also showed incoherent characteristics, with, more surprisingly, no
strong thickness dependence [26,27]. The reason
for the lack of dynamical effects with increasing thickness was because high-angle scattering
is dominated by highly localized 1s type Bloch
states, and interference with other Bloch states
is therefore suppressed. The high-angle detector
acted as a quantum state filter. Multislice image
simulations were developed based on the frozen
phonon method [28], and also predicted almost
perfect incoherent image characteristics in
crystals [29].
The combination of a simply interpretable
Figure 1
(a) Annular dark field image of diamond in the [001] orientation showing weak beam contrast from
platelet defects. (b) EEL spectra recorded on (solid line) and off (dashed line) a platelet, and the difference
spectrum showing the presence of N. Reproduced from Ref. 19 with permission.
MicroscopyandAnalysis | 25th Anniversary issue September 2012
59
scanning transmission electron microscopy
image with strong Z-contrast uncovered many
surprises in the field of interface physics. Rather
than fitting image simulations to likely structure
models, in many cases it was possible to just look
at the image and directly determine the likely
structure. Figure 2 shows a photograph of the
microscope at ORNL (note the Polaroid camera
used to record images) with an interface between
CoSi2 and Si showing an unexpected 2x1 interface
reconstruction. For reviews of some of these early
applications see [30-32]. There was also a push
to higher resolution, sub-2 Å resolution being
achieved with John Silcox’s VG microscope at
Cornell University [33, 34] and the splitting of
the dumbbell in Si<110> with the Cowley VG
machine [35]. However, a more attractive route to
higher resolution was to use a higher accelerating
voltage, as common in TEMs of the day, and VG
Microscopes produced four 300 kV STEMs, the
HB603 series, one of which, equipped with a high
resolution pole piece, was delivered to ORNL in
1993. It resolved the dumbbells in Si and could
also distinguish the sublattice in compound
semiconductors such as GaAs [36, 37].
Meanwhile, major developments had also
been taking place in the technology for EELS, in
particular the inefficient serial EELS detectors
were being replaced with vastly more efficient
parallel detectors [38]. It was natural, therefore,
to try out the possibility of achieving EELS with
atomic resolution. An efficient charge-coupleddevice parallel EELS system was acquired for
the ORNL HB501UX microscope from Dennis
McMullan, based on a design he produced for the
Cavendish VG machine [39]. Despite the small
current available in an atomic-sized beam, spectra
could be obtained simultaneously with the
HAADF signal. This allowed EELS data to be collected plane by plane across an interface using the
HAADF signal as a monitor of beam position. To
test the possibility of atomic resolution a CoSi2/
Si(111) interface was chosen, as such interfaces
were commonly atomically abrupt, and indeed
the Co L edge disappeared on moving a distance of
2.7 Å from the last plane of the silicide to the first
plane in the Si [40]. Shortly after, atomic column
resolved EELS was demonstrated at the Si/SiO2
interface [41] and sub-nm resolution mapping of
EELS fine structure was achieved [42]. 1993 was a
breakthrough year for EELS.
In 1996 VG Microscopes ceased trading, and the
increasing demand for STEM capabilities had to
be met by the traditional TEM manufacturers.
This forced them to bring their STEM performance closer to the theoretical limits [43, 44]. The
availability of an efficient STEM mode on commercial TEM columns, complete with full service
support, encouraged increasing applications and
interest in STEM. For example, the incoherent
characteristics and simple mass/thickness image
contrast made the HAADF signal ideal as the
basis for tomographic reconstruction of threedimensional (3D) structures [45, 46].
Aberration correction
Incoherent imaging has double the information limit of coherent imaging, (the cutoff being
the objective aperture diameter for incoherent
imaging, but the radius for axial phase contrast
imaging). However, the triangular nature of
the incoherent transfer function means that
the contrast is quite weak at the higher spatial
60
Figure 2
The HB501UX microscope at ORNL with
(inset) a Z-contrast
(HAADF) image of an
interface between CoSi2
and Si showing an
unexpected 2x1 reconstruction.
Figure 3
Simultaneously acquired atomic resolution EELS and EDS maps from a LaFeO3/SrTiO3 interface. Principal
component analysis was used to remove random noise. Data acquired in the JEOL JEM-ARM200F at 200
kV, courtesy of E. Okunishi (JEOL) and M. Varela (ORNL). Sample courtesy of Jacobo Santamaria’s group
(Complutense University, Spain).
frequencies. Together with the ever-present noise
problems, Z-contrast images tended to be visually
uninspiring. This changed dramatically with the
successful correction of lens aberrations [47]. Not
only could higher resolutions be achieved, but
now the same current could be put into a smaller
probe, increasing the image contrast and at the
same time improving the signal to noise ratio.
The 300 kV STEM at ORNL was fitted with a
Nion quadrupole/octupole aberration corrector
and achieved the direct imaging of the dumbbell
in Si<112> at 0.78 Å [48]. For the first time in history, the STEM was obtaining resolutions higher
than the TEM, which is in accord with physics
[49, 50] but had previously been impossible due to
the noise limitations.
EELS saw similar benefits, the first spectroscop-
ic identification of a single atom was achieved
with the ORNL HB501UX with a Nion aberration
corrector, using the original McMullan parallel
detection system [51]. Applications continued
to grow in semiconductors, ceramics, complex
oxides, catalysis and nanomaterials (for some
reviews see [52-54]).
New imaging modes also became possible. For
example, aberration correction allowed the objective lens aperture to be opened up, increasing the
resolution, but also reducing the depth of field.
This could be put to good use, focusing inside
materials to locate objects in the third dimension,
for example, locating individual stray Hf atoms
in the gate dielectric of a semiconductor device
structure [55] and the imaging of point defect
configurations inside a Si crystal [56]. While the
25th Anniversary issue September 2012 | MicroscopyandAnalysis
scanning transmission electron microscopy
lateral resolution was of atomic dimensions,
the depth resolution was much greater than for
conventional tomography, especially for objects
larger than single atoms [57]. A confocal mode of
operation overcomes this problem [58], however,
it does require two aberration correctors to be
accurately aligned with one another.
Bright field phase contrast imaging in the
STEM was also much improved. Before aberration
correction the collector aperture had to be very
small to ensure coherence, meaning the images
were noisy, however, after aberration correction
a much larger aperture could be used, collecting a
much greater fraction of the incident beam, and
the noise problem was effectively overcome. Now,
simultaneous, aberration corrected, HAADF and
phase contrast images could be compared with
pixel to pixel correlation [59]. We have also seen
the appearance of a largely incoherent bright
field phase-contrast mode, annular bright field
imaging, in which both heavy and light atoms
can be visible simultaneously [60,61]. This mode,
originally advanced by Harald Rose [62], allows
even the imaging of H atom columns in crystals
[63,64]. We have also seen the reappearance of
segmented detectors allowing differential phase
contrast images to be obtained at atomic resolution [65].
The initial 3rd order aberration correctors
were rapidly surpassed by correctors capable of
compensating up to 5th order aberrations, giving
another leap in both resolution and signal to noise
ratio. Resolution advanced first to 0.63 Å [66, 67]
and then to 0.47 Å [68, 69]. EELS also benefitted
significantly since now larger currents could
be achieved while maintaining atomic resolution, meaning that two-dimensional maps were
possible where before only line scans had been
feasible [70-72]. Even X-ray mapping has achieved
atomic resolution [73-75], and although it tends to
be more noisy than EELS due to the lower signal
collection efficiency, it is typically generated
through a more localized interaction and can
therefore give better resolution. Figure 3 compares simultaneously acquired atomic resolution
EELS and energy dispersive X-ray spectroscopy
(EDS) maps from a LaFeO3/SrTiO3 interface.
Theory has also had advances, notably, the ability to simulate core-loss images [76] as well as images formed from elastic and thermal diffuse scattering. The infamous Stobbs factor [77] was shown
not to be present for HAADF STEM images [78],
allowing more quantitative interpretation of
experimental images through simulation. EELS
fine structure imaging needs more development
however, since at present the fine structure calculations using density functional theory generally
assume a plane wave incident beam, a far cry from
an incident coherent probe. Better calculations
will need to utilize the mixed dynamic form factor to include interference terms [79].
A new world awaits exploration
One clear trend following the success of aberration correction is the use of lower accelerating
voltages to avoid knock-on damage, particularly
for graphene [80,81], not only in STEM but also in
TEM [82-84]. Since monolayer materials have little diffraction contrast there is also an advantage
in reducing the inner angle of the annular detector to improve signal to noise ratio, which has become known as medium- angle annular dark field
Figure 4
Atomic structure of a point defect complex in monolayer graphene.
(a) ADF image showing the presence of two dopant atoms.
(b) Simulated ADF image.
(c) Corresponding intensity profiles along X-X’ and Y-Y’ in the experimental and simulated images
respectively. Experimental images have been low-pass filtered.
(d-f) Si-L-, C-K-, and N-K- edge maps with blue and green squares indicating the positions of the
substitutional Si and N atoms respectively.
Scale bars are 0.2 nm. Reproduced from Ref. 88 with permission.
Figure 5
Z-contrast image of LiFePO4 showing a higher than expected intensity at some Li sites (red circle). EELS
from such sites reveals the presence of Fe, and its L2,3 ratio is higher than that of Fe in the bulk lattice sites.
Scale bar is 0.5 nm. Reproduced from Ref. 89 with permission.
MicroscopyandAnalysis | 25th Anniversary issue September 2012
61
scanning transmission electron microscopy
(MAADF) imaging although it is similar to the
original Crewe geometry. Sufficient Z-contrast
remains to directly distinguish neighboring light
elements, B, C, N and O in monolayer BN for
example [85, 86], which is a major advantage over
TEM. Atom-by-atom spectroscopy is also possible
[87, 88], and an example is shown in Figure 4. Not
only are individual atoms identified spectroscopically, they are seen to be substitutional since the
carbon elemental map shows an absence of carbon
in those regions.
Point defects can now even be probed spectroscopically. Figure 5 shows an example of an
antisite defect in LiFePO4. It shows brighter than
expected if it contained only Li, and placing the
beam on the site an EELS spectrum can be recorded with sufficient signal not only to identify
the presence of Fe but to determine its valence.
Surprisingly, the valence of the Fe on the Li site
is higher than Fe in the bulk lattice, which can be
explained by the trapping of mobile Li vacancies
[89]. The ability to perform mapping based on
EELS fine structure allows not only valence to be
mapped, but also new properties. Figure 6 shows
how spin states can be mapped in La0.5Sr0.5CoO3-d
thin films to reveal the presence of a spin state
superlattice, based on changes in the pre-peak
on the O K edge [90]. Recently, the localized Cu
3d hole in the bc plane of La2CuSnO6 was imaged
using fine structure on the O K edge [91].
We are also beginning to realize that not all
beam-induced changes to the specimen are bad.
A gentle provocation can cause the specimen
to explore metastable configurations [92, 93].
Density functional theory can then be used to
reveal diffusion pathways and other properties.
An example is presented in Figure 7, which shows
sub-2 nm CdSe nanoclusters that are known to
emit white light. This white light emission is
remarkable because it occurs even from single
nanoclusters, meaning that it cannot simply be
explained by an ensemble of clusters with different emission characteristics. Observation in the
STEM revealed a continuously changing atomic
structure, presumably induced by the beam. We
realized that excitation by an ultraviolet photon
might also impart enough energy to the cluster
to induce a fluxional state. Quantum mechanical
molecular dynamics confirmed the conjecture,
showing how the fluxionality would lead to a
continuously changing band gap, hence the white
light emission [94].
These new tools are opening windows onto
many areas previously inaccessible. In catalysis it
is possible to see the structure of the smallest clusters, and through intensity analysis determine
their 3D structure and alloy composition (see
Figure 8 [95]). In thermoelectric materials it is possible to see vacancies and how they affect thermal
vibration amplitudes [96]. The atomic structures
of dopant nanoclusters can be determined within
semiconductors [97] and single Au atoms have
recently been seen in GaAs nanowires [98]. This is
just a glimpse of the remarkable new insights now
being gained with these techniques into both new
and more traditional materials.
Future directions
There are several new directions currently being
pursued that will present even more novel opportunities in the future. For example, the ability
to obtain probes of atomic dimensions carrying
62
Figure 6
Spin state mapping in La0.5Sr0.5CoO3-d thin films. The Z-contrast image shows a superlattice induced by
vacancy ordering, the Co plane containing vacancies indicated by the blue arrow, the stoichiometric
plane by the red arrow. The O K EELS edge (left) is different from the two planes, indicating a spin state
superlattice. Density functional calculations (right) show the strong spin polarization in the blue planes
near the Fermi level (dashed line). Reproduced from Ref. 90 with permission.
Figure 7
(a,b) Z-contrast images of a sub-2 nm CdSe nanocluster showing its fluxional state. (c) White light emission
spectrum from sub-2 nm nanoclusters. (d) Density functional calculations reveal a continuously varying
band gap. Spectrum colour represents the energy gap. Reproduced from Ref. 94 with permission.
orbital angular momentum offers new possibilities to measure magnetic properties [99]. There is
much effort in building in-situ stages that maintain atomic resolution [100] (see for example the
ordering of ultrasmall FePt nanoparticles [101]).
Furthermore, it is now possible to achieve atomic
resolution at one atmosphere [102] and even in
liquids [103, 104].
On the instrumental front we see the potential
for chromatic aberration correction in STEM,
especially at the lower accelerating voltages now
in demand where chromatic aberration becomes
the limiting factor for spatial resolution [105]. A
number of monochromator designs are appearing
that would be very appealing if incorporated
into an aberration-corrected STEM [105-108],
providing the ability to map band gaps and the
plasmonic response of nanomaterials [109, 110] .
Cathodoluminescence mapping on the STEM is
another way to probe optical properties, complementary to EELS, although generally expected to
show poorer spatial resolution [111, 112].
It is truly amazing how the STEM has improved
in the last 25 years, changing from a dedicated
microscope for sub-nanometer analysis to an
instrument capable of single atom imaging and
spectroscopy with sub-Ångstrom resolution.
Surely there can be no better way to unravel the
complex functionalities of nanomaterials – every
particle is likely to have different structure, different bonding, different impurities and different
properties. Examination atom-by-atom, bond-bybond, in conjunction with local measurements
of properties and interpreted through theoretical
calculations now offers a new road to understanding the ultimate atomic origins of materials
properties.
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25th Anniversary issue September 2012 | MicroscopyandAnalysis
scanning transmission electron microscopy
biography
Stephen J. Pennycook
is a Corporate Fellow
at Oak Ridge National
Laboratory and a fellow
of five societies including
the Materials Research
Society, which awarded
him their 2012 Innovation
in Materials Characterization Award.
Figure 8
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MicroscopyandAnalysis | 25th Anniversary issue September 2012
abstract
Scanning transmission electron
microscopy has undergone dramatic
changes during the last 25 years, moving
from a niche activity by just a few
groups with dedicated microscopes to
the mainstream technique for atomic
resolution imaging and analysis that it
is today. This article tracks that progress,
beginning with a glimpse of the state of
the art 25 years ago, highlighting the key
advances in the field, and ending with
some examples of today’s state of the art.
acknowledgements
The author would like to acknowledge
fruitful interactions with numerous
colleagues over many years, in particular
L. M. Brown, A. Howie, N. D. Browning, M.
F. Chisholm, P. D. Nellist, A. R. Lupini, M.
Varela, A. Y. Borisevich, O. L. Krivanek, M.
P. Oxley, S. T. Pantelides and J-C. Idrobo,
and support from the US Department of
Energy, Office of Basic Energy Sciences,
Materials Sciences and Engineering
Division.
corresponding author details
Dr Stephen Pennycook
Scanning Transmission Electron
Microscopy Group
Materials Science and Technology Division
Oak Ridge National Laboratory,
PO Box 2008, Oak Ridge, TN 37831-6071,
USA
Tel: +1 865 574 5504
Email: [email protected]
Microscopy and Analysis 26(6):59-64 (AM),
2012
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25th Anniversary issue September 2012 | MicroscopyandAnalysis