Strain Localization in Thin Films of Bi(Fe,Mn)O3 Due to the Formation of Stepped Mn(4+)-Rich Antiphase Boundaries

Abstract

The atomic structure and chemistry of thin films of Bi(Fe,Mn)O3 (BFMO) films with a target composition of Bi2FeMnO6 on SrTiO3 are studied using scanning transmission electron microscopy imaging and electron energy loss spectroscopy. It is shown that Mn(4+)-rich antiphase boundaries are locally nucleated right at the film substrate and then form stepped structures that are approximately pyramidal in three dimensions. These have the effect of confining the material below the pyramids in a highly strained state with an out-of-plane lattice parameter close to 4.1 Å. Outside the area enclosed by the antiphase boundaries, the out-of-plane lattice parameter is much closer to bulk values for BFMO. This suggests that to improve the crystallographic perfection of the films whilst retaining the strain state through as much of the film as possible, ways need to be found to prevent nucleation of the antiphase boundaries. Since the antiphase boundaries seem to form from the interaction of Mn with the Ti in the substrate, one route to perform this would be to grow a thin buffer layer of pure BiFeO3 on the SrTiO3 substrate to minimise any Mn-Ti interactions.

Full text

N A N O E X P R E S S Open Access
Strain Localization in Thin Films of Bi(Fe,Mn)O
3
Due to the Formation of Stepped Mn
4+
-Rich
Antiphase Boundaries
I MacLaren
1*
,BSala
1
, S M L Andersson
1
, T J Pennycook
2,3,7
, J Xiong
4,5
,QXJia
4
, E-M Choi
6
and J L MacManus-Driscoll
6
Abstract
The atomic structure and chemistry of thin films of Bi(Fe,Mn)O
3
(BFMO) films with a target composition of
Bi
2
FeMnO
6
on SrTiO
3
are studied using scanning transmission electron microscopy imaging and electron energy
loss spectroscopy. It is shown that Mn
4+
-rich antiphase boundaries are locally nucleated right at the film substrate
and then form stepped structures that are approximately pyramidal in three dimensions. These have the effect of
confining the material below the pyramids in a highly strained state with an out-of-plane lattice parameter close to
4.1 Å. Outside the area enclosed by the antiphase boundaries, the out-of-plane lattice parameter is much closer to
bulk values for BFMO. This suggests that to improve the crystallographic perfection of the films whilst retaining the
strain state through as much of the film as possible, ways need to be found to prevent nucleation of the antiphase
boundaries. Since the antiphase boundaries seem to form from the interaction of Mn with the Ti in the substrate,
one route to perform this would be to grow a thin buffer layer of pure BiFeO
3
on the SrTiO
3
substrate to minimise
any Mn-Ti interactions.
Keywords: Bismuth ferrite; Scanning transmission electron microscopy (STEM); Strain; Thin films; Multiferroic;
Antiphase boundaries
Background
There has been considerable interest in bismuth ferrite
due to the fact that it supports simultaneous permanent
magnetic [1] and ferroelectric [2] orderings, has poten-
tial for use as a multiferroic with strong magnetoelectric
couplings [3, 4], and is a rare low bandgap ferroelectric,
which is tunable across the whole visible spectrum by
B-site doping, and hence could be used as an absorber
material with enhanced carrier extraction [5, 6]. Neverthe-
less, it has proved difficult to create the desired properties
in pure BiFeO
3
, and attention has turned to using com-
positional modification [7–10] and/or substrate-induced
strain [11, 12] to alter the structure and properties of the
material. Recently, Choi et al. [13] reported the growth of
ferroelectric and ferromagnetic films of nominal compos-
ition of BiFe
0.5
Mn
0.5
O
3
on SrTiO
3
and showed that very
careful slow growth was critical to achieving this result. In
this work, we report a careful atomic resolution investiga-
tion of antiphase boundaries forming close to the film-
substrate interface and show, using a careful study with
atomic resolution scanning transmission electron mi-
croscopy, that even in such high-quality films, there are
additional, hitherto, unexpected complexities driven by a
tendency to some chemical segregation. These result in
the formation of additional antiphase boundaries that have
the effect of releasing the compressive elastic strain
fromtheepitaxialgrowthonthesubstrateabovethe
boundaries.
Methods
Films were grown by pulsed laser deposition onto a SrTiO
3
(STO) substrate held at 640 °C from a Bi(Fe,Mn)O
3
(BFMO) ceramic target of composition Bi
2
FeMnO
6
using a
KrF laser pulsed at 2–5 Hz and an oxygen pressure of
100 mTorr, as described in more detail in our previous
publication [13]. Samples were prepared for scanning
transmission electron microscopy using a modified focused
ion beam liftout procedure. Firstly, the sample was coated
* Correspondence: ian.maclaren@glasgow.ac.uk
1
SUPA School of Physics and Astronomy, University of Glasgow, Glasgow
G12 8QQ, UK
Full list of author information is available at the end of the article
© 2015 MacLaren et al. Open Access This article is distributed under the terms of the Creative Commons Attribution 4.0
International License (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use, distribution, and
reproduction in any medium, provided you give appropriate credit to the original author(s) and the source, provide a link to
the Creative Commons license, and indicate if changes were made.
MacLaren et al. Nanoscale Research Letters (2015) 10:407
DOI 10.1186/s11671-015-1116-8

with about 30 nm of carbon; the samples were then pre-
pared using a standard focused ion beam (FIB) liftout pro-
cedure, with the orientation of the slice chosen to ensure
that the section had a primitive {001} plane of the perovsk-
ite in the sample plane. After liftout, the sample was at-
tached onto the side of an Omniprobe copper mount using
platinum deposition and then thinned to ~100 nm using
5-kV Ga ions. Final thinning was then performed using a
Gatan PIPS equipped with the low-energy upgrade, and
the Ar ions were directed onto the sample using angles
of +8° and −8° with single-sided sector milling engaged
and the orientation of the sample set to ensure that no
copper was re-deposited from the support onto the lifted
out slice. This final thinning was performed at a low volt-
age of 500 V for about half an hour.
Scanning transmission electron microscopy was per-
formed using two microscopes. High-angle annular
dark-field (HAADF) imaging was performed using a
JEOL ARM200F with a cold field emission source and
using an accelerating voltage of 200 kV, a probe of semi-
convergence angle of 29 mrad and an effective collection
angle range for the HAADF detector of about 90–
150 mrad. Images were collected at short exposure times
of about 10 μs/pixel using a mode of repeated scanning
to record 10–30 frames of the same structure for post-
processing.
Electron energy loss spectroscopy-spectrum imaging
(EELS-SI) was performed using a NION UltraSTEM op-
erated at 100 kV and a Gatan Enfina EELS with a probe
angle of 32 mrad and a collection angle of 37 mrad.
Initial processing for quantitative analysis of atom
positions was performed by sorting out a number of
images from the image stack that are all free of obvious
glitches and distortions. These were then aligned using
the SDSD plug-in for Digital Micrograph [14] and
summed to create a high signal to noise, drift-free image,
as done in our previous work [15–18]. Quantitative de-
termination of the atomic positions in the thin film was
performed using an experimental Image Analysis plug-in
for Gatan Digital Micrograph (courtesy of Dr Bernhard
Schaffer, Gatan GmbH), which automatically detects the
peaks and fits them using a 2D Gaussian function and
then provides a list of peak positions and fit parameters.
Further processing, evaluation and plotting of the peak
list were then performed using conventional spread-
sheets, detailed in previous publications [15–18].
EELS-SI datasets were processed by the following pro-
cedure. The datasets were first processed to remove any
extraneous X-ray spikes resulting from X-ray generation
from stray scattering in the spectrometer. Multivariate
statistical analysis was then performed using the plug-in
of Lucas et al.[19] to separate the real elemental edges
from the random noise. The quantification was then
performed using an EELS modelling approach [20–23]
included in an experimental Digital Micrograph plug-in
[24]; since the low-loss datasets were not available, only
relative quantification of atomic ratios was performed
and the spectra were fitted from 400 to 760 eV to cover
the Ti-L
2,3
, O-K, Mn-L
2,3
and Fe-L
2,3
edges.
Valence state mapping was performed using multiple
linear least squares (MLLS) fitting of the region covering
the L
3
and L
2
white lines for Mn using internal references
from the datasets for the Mn
3+
and Mn
4+
near edge
shapes, on the assumption that the whole film exists as
some mixture of these two oxidation states (in accordance
with our previous work [13]).
Results
Figure 1 shows a typical HAADF STEM image of an
area of the BFMO film containing antiphase boundaries.
These may be recognised as the same stepped structure
of antiphase boundary as has recently been reported by
MacLaren et al.in Bi
0.85
Nd
0.15
Fe
0.9
Ti
0.1
O
3
[18]. It is clear
that these stepped structures are nucleated at or just
above the interface with the SrTiO
3
and then grow step-
wise upwards until they intersect another stepped
boundary nucleated elsewhere at the substrate-film inter-
face, creating a pyramidal structure. There is also evidence
from overlapped regions, such as that indicated in Fig. 1,
and from focal series collected in some areas, that the
overall 3D structure is of pyramids of BFMO connected in
direct epitaxial relationship to the SrTiO
3
separated from
the outer part of the film by this pyramidal antiphase
boundary (APB) network. Thus, the A-site positions in
the outer part of the film are always in an antiphase rela-
tionship with the A-sites in the SrTiO
3
substrate: this has
Fig. 1 HAADF image of the BFMO thin film on SrTiO
3
.Two lines are
shown for profiles taken
MacLaren et al. Nanoscale Research Letters (2015) 10:407 Page 2 of 7

been observed in multiple TEM samples and for two dif-
ferent film thicknesses.
Figure 2a shows a plot of the out-of-plane lattice par-
ameter along the pink line shown in Fig. 1, which passes
through the location where the APB is flat and about 1
unit cell above the STO-BFMO interface. This shows a
jump to about 4.15 Å at the film-substrate interface,
followed by larger jumps to either side of the APB just
above the interface, with a short spacing of about 3.1 Å
in between. This is entirely in accord with the previously
published structure of flat terraces on such APBs [16],
where the first cell on either side of the APBs has a huge
cparameter of about 4.3–4.4 Å due to the stabilisation
of a super-tetragonal, highly polar phase in response to
the high charge density at the boundary [16]. Above the
boundary, the out-of-plane parameter decays gradually
over 5 or 6 unit cells back to an equilibrium level,
which is slightly higher than that in the SrTiO
3
.
Using the average spacing in the SrTiO
3
as an in-
ternal calibration of 3.905 Å, this new level is a little
higher at ~3.93 Å. For comparison, the in-plane lattice
parameter is plotted in Fig. 2b along the same line. In this
case, this is almost the same in the BFMO as in the STO,
with just a slight disturbance by the internal structure of
the APB.
Figure 2c shows a plot of the out-of-plane lattice par-
ameter along the green line in Fig. 1. This line runs
through the centre of a pyramid, and the plot is surpris-
ingly different. Specifically, there is no lattice parameter
jump at the film-substrate interface, but the whole area
inside the pyramid has a large out-of-plane parameter of
about 4.08 Å. There is then the usual big peak to
about 4.4 Å, due to the stabilisation of a super-
tetragonal polar-ordered phase by the electric field from
the boundary [16]; the drop to ~3 Å; and the second peak
of about 4.4 Å (as explained above), when the line passed
through the APB at the tip of the pyramid. This is
followed by the decay back to an equilibrium value of the
lattice parameter, which in this case is ~3.94 Å. Again, for
comparison, the in-plane parameter is shown for compari-
son in Fig. 2d. As for Fig. 2b, there is no significant change
of the in-plane parameter from the STO value, either
inside or outside the pyramid, except the minor distur-
bances at the APB itself.
Fig. 2 Quantification of the structure in Fig. 1. aOut-of-plane and bin-plane lattice parameters as a function of distance along the pink line
shown in Fig. 1. cOut-of-plane and din-plane lattice parameters as a function of distance along the green line shown in Fig. 1. In all cases, the
position values <0 are in the SrTiO
3
substrate, and position values >0 are in the BFMO film
MacLaren et al. Nanoscale Research Letters (2015) 10:407 Page 3 of 7

All this provides an interesting counterpart to previ-
ously published work where crystallographic analysis
was mainly performed using X-ray diffraction [13], pro-
viding highly accurate average values of film and sub-
strate lattice parameters, but little idea of the detailed
variation of the crystal parameters within the film. The
present work clearly supports the conclusions of our
previous publication that the films are highly coherent
with the substrate and that the in-plane parameters are
constrained to those of the STO. However, it is now
clear that the out-of-plane parameter shows locally sig-
nificant variations within the film. Those areas that are
in direct epitaxial contact with the substrate show a sig-
nificantly enhanced c:aratio with cvalues close to 4.1 Å.
This is significantly above the total cparameter for the
films from X-ray diffraction of about 4.015 Å [13]. This
is compensated for by the fact that outside the APB, the
cparameter decays back to something not far from the
expected bulk value. This range of cparameters will give
an average for the whole film not far from 4 Å and a ra-
ther spread reciprocal space peak for the film, as ob-
served in Fig. 1c of Choi et al.[13].
Figure 3 shows an atomic resolution chemical map of
the area where an APB is in direct contact with the
underlying SrTiO
3
. It should be noted that the APB does
not lie right on the interface, even on the left-hand side,
but one cell inside the BFMO. Chemical maps show that
there is an enrichment of Mn on the B-sites at the inter-
face on the left-hand side. This corresponds to where
the APB runs right along the film-substrate interface,
and there is also a small but significant concentration of
Ti along with the Mn on the B-sites at the interface, as
reported previously by Choi et al.[13]. On the right-
hand side, where the APB steps up and away from the
film-substrate interface, there is also a strong enrich-
ment of Mn to the boundaries. The fact that Mn is
already segregating to the film-substrate interface and
that a flat APB forms here clearly shows that Mn is asso-
ciated with the nucleation of the APBs in the first unit
cell above the film-substrate interface, possibly in associ-
ation with a little Ti from the substrate. In our previous
studies [16], it was found that these boundaries were
formed in BiFeO
3
in the presence of excess Ti
4+
,andit
seems that Mn can play a similar role. It is then seen
that the Mn segregation is key to how these boundaries
then step away from the film-substrate interface into the
pyramid structures shown in Fig. 1.
Figure 4 shows atomic resolution chemical maps of a
stepped area of the APB showing that, whilst Fe is
present both in B-sites in the surrounding perovskite
and in the boundary, Mn is strongly segregated to the B-
sites in the boundary. It was clear from studying the EEL
spectra from this spectrum image, as shown in Fig. 3d,
that the Mn white lines were shifting and changing
shape between the perovskite and the boundary, suggest-
ing that the oxidation state is changing. The trends are
much like those seen previously by Garvie and Craven
[25], whereby the edge onset moves to higher energy
with increasing oxidation state, with a concurrent de-
crease in the L
3
:L
2
ratio, as later used by Wang et al. and
others [26–28] for measuring Mn oxidation states. Map-
ping using the L
3
:L
2
ratio was used by Choi et al.[13] to
show a change of Mn oxidation state in these films from
Mn
3+
in the perovskite to Mn
4+
on the interface to the
Fig. 3 Nucleation of an antiphase boundary at the SrTiO
3
:BFMO interface. (left) HAADF image showing antiphase boundary features either at
(left) or above (right) the film-substrate interface, as well as showing the area used for EELS-SI. (right) Elemental map from processing of the EELS-
SI data where Fe is red,Mnisgreen,Tiisblue and the simultaneously acquired HAADF signal is purple
MacLaren et al. Nanoscale Research Letters (2015) 10:407 Page 4 of 7

SrTiO
3
. In the current work, MLLS fitting was used to
treat a background-subtracted Mn edge spectrum image
of this area (640–665 eV energy loss) as a linear sum of
two extreme components, one from the matrix and an-
other from the step in the boundary. The former must
be Mn
3+
in accord with Choi et al.[13], and the latter
was believed to be close to Mn
4+
. The results are shown
in Fig. 3c which clearly show that all the Mn in the
boundary steps oxidises to Mn
4+
.
Discussion
With the benefit of this atomic resolution chemical ana-
lysis, it is now possible to compare this stepped APB
structure to that previously reported in Nd,Ti co-doped
BiFeO
3
[18]. In that work, flat APB sections were formed
with Ti
4+
at the core, and the steps on APBs were always
formed from four closely separated columns of Fe
3+
,
which are linked together as edge-sharing FeO
6
octahe-
dra. In this work, it is shown that very similar steps can
be created in a BiFeO
3
derivative (Bi(Fe,Mn)O
3
) using
Mn
4+
to form the steps, in preference to Fe
3+
. One rea-
son for the segregation of Mn to these boundaries could
be that the desired BiFe
0.5
Mn
0.5
O
3
composition is less
stable than a mixture of Fe-rich Bi(Fe,Mn)O
3
and phases
richer in Mn. If this were the case, there would be a
clear driving force to phase separation, which when it
occurs for slow film deposition rates would lead mainly
to the formation of local, non-stoichiometric defects, ra-
ther than larger scale phase separation into second
phases. The fact that there is very little Ti present in the
film, apart from possibly a little diffused from the sub-
strate in the first atomic layer, may explain why flat APB
sections are only supported right at the interface with
the substrate, seeing as previous observations found Ti
was vital to the stabilisation of flat APBs [16].
There is then the question of why the Mn changes oxi-
dation state to Mn
4+
in the boundaries. It is well known
that these boundaries contain an excess of negative
charge in the core [16, 18]: this conclusion was easily
reached in our previous publications by counting oxygen
ions and the cations and calculating the net charge dens-
ity. One way that this excess charge could be reduced, if
not eliminated entirely, would be to increase the oxida-
tion state of the cations. This, therefore, explains why
Mn in the boundaries is in the higher Mn
4+
oxidation
state and may provide an additional energetic reason
why Mn is preferred to Fe for the cores of the steps,
since Mn readily supports higher oxidation states than
Fe. Whatever is the case, the stepped boundaries are still
negatively charged in the core, and this is estimated to
correspond to a charge density of −0.55 C m
−2
for the
core as Mn
4+
, as compared to the previously published
figure of −1.09 C m
−2
for the core as Fe
3+
[18]. The fact
that the boundary remains charged is also attested to by
Fig. 4 Atomic resolution chemical maps of a stepped region of an antiphase boundary. aSurvey image. bComposite image of Fe (red), HAADF
signal (purple) and Mn signal (green). cComposite image showing the Mn
3+
(lilac) and Mn
4+
(pink) MLLS fits. dStandard spectra used for this
MLLS fit
MacLaren et al. Nanoscale Research Letters (2015) 10:407 Page 5 of 7

the fact that the surrounding material is polarised to-
wards the APB with the Bi atoms close to the APB not
sitting in the middle of the space between four sur-
rounding B-sites but always offset towards the APB, as
also observed in our previous studies [16, 18].
Since it is shown that the strain in the film is maxi-
mised when the film is in direct epitaxial contact with
the substrate, and relaxed by the presence of Mn-rich
APBs, it is clear what must be done to further improve
the properties of these materials. APBs are nucleated at
the film-substrate interface, and this seems to be associ-
ated with the preferential segregation of Mn
4+
to the B-
sites next to the interface, which possibly occurs due to
the interaction with Ti
4+
ions diffusing into the first layer
or two of the film from the substrate. It would therefore
seem sensible to suggest that future developments in
film growth should concentrate on the elimination of
Mn diffusion to the film-substrate interface. This could,
perhaps, be achieved by first depositing a few atomic
layers of pure BiFeO
3
before turning the deposition to
BiFe
0.5
Mn
0.5
O
3
to prevent any interaction of Mn and Ti.
Conclusions
In conclusion, using atomic-resolution structural and
spectroscopic characterisation we have found that the
growth of bismuth ferrite manganite of target compos-
ition BiFe
0.5
Mn
0.5
O
3
on SrTiO
3
results in the formation
of Mn
4+
-rich APBs emanating from the film-substrate
interface. These then form a stepped structure and inter-
sect with one another to form a structure of approxi-
mate pyramids of film in direct 1:1 epitaxial relationship
to the substrate. The effect of the pyramidal APBs is to
leave the outer layer of the film in an antiphase relation-
ship to the substrate. This has the effect of concentrating
strain close to the interface but relaxing it in the outer
layers of the film. The formation of the APBs appears to
be driven by a tendency towards phase separation, namely
into Mn-rich and Mn-poor phases.
Finally, more broadly, the finding of this work high-
lights the challenges involved in depositing homoge-
neous complex oxide thin films on the limited number
of mostly lattice mismatched single-crystal substrates
which are available and presents the case for use of suit-
able buffer layers to prevent deleterious chemical interac-
tions in film growth.
Competing interests
The authors declare that they have no competing interests.
Authors’contributions
QXJ and JX grew the film, EMC performed much of the basic film characterisation,
TJP performed the atomic resolution spectroscopy, IM performed the atomic
resolution imaging and oversaw the electron microscopy data processing, BS
performed the majority of the electron microscopy data processing and image
preparation, SMLA assisted with the processing of oxidation state standards and
fitting the experimental data to standard spectra for oxidation state mapping, JLD
designed and oversaw the research programme. All authors contributed to
scientific discussions and the preparation of the manuscript.
Acknowledgements
This work was supported by the European Research Council (ERC-2009-AdG
247276 NOVOX). The work at Los Alamos National Laboratory was supported
by the LDRD programme and was performed at the Center for Integrated
Nanotechnologies (CINT), an Office of Science User Facility operated for the
US Department of Energy (DOE) Office of Science. The authors are indebted
to the continuing support of the EPSRC for the SuperSTEM facility, which
made this work possible. We are also grateful to SUPA and the University of
Glasgow for the funding of the JEOL ARM200F, which was also used in this
work. The assistance of Mr William (Billy) Smith in the preparation of the FIB
specimen is gratefully acknowledged.
Author details
1
SUPA School of Physics and Astronomy, University of Glasgow, Glasgow
G12 8QQ, UK.
2
SuperSTEM Laboratory, STFC Daresbury Laboratories, Keckwick
Lane, Warrington WA4 4AD, UK.
3
Department of Materials, University of
Oxford, Parks Road, Oxford OX1 3PH, UK.
4
Center for Integrated
Nanotechnologies, Los Alamos National Laboratory, Los Alamos, NM 87545,
USA.
5
State Key Lab of Electronic Thin Films and Integrated Devices,
University of Electronic Science and Technology of China, NO.4, Section 2,
North Jianshe Road, Chengdu 610054, China.
6
Department of Materials
Science, University of Cambridge, 27 Charles Babbage Road, Cambridge CB3
0FS, UK.
7
Present Address: Physics of Nanostructured Materials, Faculty of
Physics, University of Vienna, Boltzmanngasse 5, A-1090 Vienna, Austria.
Received: 7 July 2015 Accepted: 12 October 2015
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