Physica C 355 (2001) 225±230
www.elsevier.nl/locate/physc
TEM of ultra-thin DyBa2Cu3O7±x ®lms deposited on
TiO2 terminated SrTiO3
S. Bals a,*, G. Rijnders b, D.H.A. Blank b, G. Van Tendeloo a
b
a
EMAT, University of Antwerp (RUCA), Groenenborgerlaan 171, B-2020 Antwerp, Belgium
Low Temperature Division, MESA+ Research Institute, Applied Physics, University of Twente, P.O. Box 217, 7500 AE Enschede,
The Netherlands
Received 13 November 2000; received in revised form 2 January 2001; accepted 5 January 2001
Abstract
Using pulsed laser deposition ultra-thin DyBa2 Cu3 O7±x ®lms were deposited on a single terminated (0 0 1) SrTiO3
substrate. The initial growth was studied by high-resolution electron microscopy. Two di€erent types of interface
arrangements occur and were determined as: bulk±SrO±TiO2 ±BaO±CuO±BaO±CuO2 ±Dy±CuO2 ±BaO±bulk and bulk±
SrO±TiO2 ±BaO±CuO2 ±Dy±CuO2 ±BaO±CuO±BaO±bulk. This variable growth sequence causes structural shifts, resulting in antiphase boundaries with displacement vector R ˆ ‰0 0 13Š, as well as local chemical variations. Ó 2001
Elsevier Science B.V. All rights reserved.
PACS: 68.35.Ct; 68.55.-a
Keywords: Interface arrangement; Antiphase boundary; Growth mode
1. Introduction
Although a lot of research has been dedicated
to the growth of HTSC thin ®lms on various
substrates (see Refs. [1,2]), epitaxial growth of
these complex oxide ®lms is not yet fully understood. Di€erent techniques such as thermal evaporation [3,4], molecular beam epitaxy (MBE) [5,6],
pulsed laser deposition (PLD) [7±9], etc., have
been tried in order to obtain smooth epitaxial superconducting ®lms. Experience has learned that
the initial growth of these ®lms is of extreme importance. Di€erent techniques are used to analyse
*
Corresponding author. Tel.: +32-321-80249; fax: +32-32180257.
E-mail address: [email protected] (S. Bals).
the initial growth modes and interface arrangements: using re¯ection high-energy electron diffraction (RHEED) it is possible to monitor ®lm
growth during deposition; atomic force microscopy (AFM) can be used to study the surface
morphology after deposition. Transmission electron microscopy (TEM) and particularly highresolution electron microscopy (HREM) allow to
deduce the atomic stacking at the interface. For
YBa2 Cu3 O7±x thin ®lms on (0 0 1) SrTiO3 this has
been done by various authors [4,10]. Wen et al. [4]
reported a single type of atomic stacking at the
interface: bulk±SrO±TiO2 ±BaO±CuO2 ±Y±CuO2 ±
BaO±CuO±BaO±bulk. However, X-ray standing
wave measurements [11] suggest the existence of a
variable stacking mode with an extra BaO block at
the interface, leading to the possibility of partial
overgrowth at unit cell steps at the substrate.
0921-4534/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved.
PII: S 0 9 2 1 - 4 5 3 4 ( 0 1 ) 0 0 0 3 4 - X
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S. Bals et al. / Physica C 355 (2001) 225±230
These observations, however, were made on annealed SrTiO3 substrates, which means that the
atomic surface layer of SrTiO3 can be either SrO
or TiO2 (mixed termination). In the present study
we used a single terminated (0 0 1) SrTiO3 substrate with TiO2 as surface layer. The initial
growth of a DyBa2 Cu3 O7±x thin ®lm was studied
by HREM and the layer sequence of the ®rst
DyBa2 Cu3 O7±x unit cell was determined at di€erent positions of the ®lm.
both under ideal conditions is not straightforward; the lattice mismatch between ®lm and
substrate (about 2%) causes local bending of the
ultra-thin ®lm. No interface dislocations are present (because of the reduced thickness of the ®lm)
and no surface steps in the substrate have been
noticed. Fig. 1a gives an overview of the DyBa2 Cu3 O7±x ®lm on the SrTiO3 substrate; the epitaxy
is perfect all over the ®lm with [0 0 1]DyBCO //
[ 0 0 1]STO . The most prominent defects occurring
2. Experimental
Using PLD, an ultra-thin ®lm DyBa2 Cu3 O7±x
with a thickness of approximately 6 nm was deposited on a (0 0 1) SrTiO3 substrate. The energy
source was an KrF excimer laser (248 nm). A
sintered ceramic target was used which contained
cation atoms in nominal ratio Dy:Ba:Cu ˆ 1:2:3.
During deposition, the ®lm growth was monitored
in situ by RHEED under high oxygen pressure [9].
The substrate had a miscut angle less than 0.2° and
an atomically smooth single terminated TiO2 surface, obtained as described in Ref. [12]. During
growth, the substrate was heated to a temperature
of 780°C and an oxygen background pressure of
0.18 mbar was applied. On top of the ®lm a protective SrTiO3 layer was deposited. During the
deposition of this layer, the substrate temperature
was 400°C and the oxygen background pressure
was 0.04 mbar.
Cross-section samples for TEM were cut parallel to the (0 1 0) or (0 0 1) plane of SrTiO3 . After
mechanically polishing the sample down to a
thickness of 10 lm, electron transparency was
obtained by ion milling. Electron microscopy was
performed with a Jeol 4000EX, operated at 400 kV
Image
and having a point resolution of 1.7 A.
simulations were carried out using the M A C T E M P A S software.
3. Results
HREM is performed along the perovskite cube
directions; i.e. along [1 0 0] or [0 1 0] of the DyBa2 Cu3 O7±x ®lm. Imaging the ®lm and the substrate
Fig. 1. (a) Cross-sectional overview of a DyBa2 Cu3 O7±x thin
®lm on a SrTiO3 substrate. A protective SrTiO3 layer was deposited on top of the ®lm. The interface is marked by an arrow.
(b) HREM micrograph showing an antiphase boundary with
displacement vector R ˆ ‰0 0 13Š: The boundary is perpendicular
to the interface and extends over the total thickness of the ®lm.
S. Bals et al. / Physica C 355 (2001) 225±230
(and which certainly in¯uence the properties of the
®lm) are antiphase boundaries (see Fig. 1b). They
start at the substrate±®lm interface and persist
over the total ®lm thickness. These antiphase
boundaries are perpendicular to the interface and
can be described by a displacement vector R ˆ
‰0 0 13Š. Various authors already reported the presence of such boundaries in thin REBa2 Cu3 O7±x
®lms, grown on SrTiO3 [4,10]. In these cases the
presence of the antiphase boundaries was attributed to unit cell steps (0.39 nm) at the SrTiO3
substrate. In the present study, however, the interface between ®lm and substrate is atomically
¯at and no amorphous or secondary phases were
observed.
In principle there are 12 arrangements possible
for the interface between SrTiO3 and DyBa2 Cu3 O7±x : SrTiO3 can end with SrO or TiO2 and
the DyBa2 Cu3 O7±x unit cell contains six di€erent
(0 0 1) layers. However, due to the surface treatment of the substrate, the surface atomic layer of
SrTiO3 is known to be a TiO2 layer. This limits the
number of possible interface stackings to 6.
Fig. 2 shows a detailed HREM image of the
interface for a very small sample thickness and a
defocus around )25 nm. Image simulations indicate that the cations are imaged as dark dots and
that oxygen columns are imaged as bright dots.
The substrate±®lm interface is indicated by two
arrowheads in Fig. 2; an antiphase boundary is
present in the middle of the ®gure.
Since SrTiO3 has the perovskite ABO3 structure
while DyBa2 Cu3 O7±x is closely related to perovskite, we can assume that the perovskite stacking will remain continuous. In DyBa2 Cu3 O7±x the
A sites are occupied with Dy and Ba ions, while
the Cu ions occupy the B sites. Since the substrate
has a B site (TiO2 ) termination, the ®rst layer of
the ®lm is expected to be an AO layer in order not
to disturb the perovskite stacking. This rules out
the possibility of a CuOx layer next to the TiO2
layer and limits the number of possible interface
con®gurations to three; they are represented
schematically in Fig. 3.
Intensity scanning along the vertical [0 0 1]
axis allows to determine the ®rst layer of the
DyBa2 Cu3 O7 x ®lm (see inset of ®g. 2). This ®rst
``DyBa2 Cu3 O7 x '' layer was found to be at the
227
Fig. 2. HREM micrograph of the interface region. Two
arrowheads indicate the interface. In the middle of the picture,
an antiphase boundary is present (marked by the heavy arrow).
In this area, slight bending of the atomic (0 0 1) planes, can be
seen. Line scans along the [0 0 1] direction are shown for both
sides of the antiphase boundary. Image simulation are shown as
insets. For the left side of the antiphase boundary the following
model was simulated: bulk±SrO±TiO2 ±BaO±CuO±BaO±CuO2 ±
Dy±CuO2 ±BaO±bulk (parameters used: d ˆ 2:4 nm, d ˆ 24
nm). The simulation on the right side of the antiphase boundary
is based on the interface model: bulk±SrO±TiO2 ±BaO±CuO2 ±
Dy±CuO2 ±BaO±CuO±BaO±bulk (parameters used: d ˆ 3:2
nm, d ˆ 24 nm).
same level on both sides of the antiphase boundary, excluding a step at the SrTiO3 surface as the
reason for the antiphase boundary formation.
is larger
Because the size of a Ba-block (4.14 A)
than the size of a Dy-block (3.39 A), it is fairly
easy, using the line intensity scan, to locate the
position of the Dy layer on both sides of the antiphase boundary (see Fig. 2). We observed a clear
shift of the Dy layer over R ˆ ‰0 0 13Š on both sides
of the antiphase boundary, leading to a chemical
and structural mis®t and unavoidably two di€erent
interface con®gurations. The structural mis®t and
the change in interatomic spacing at the antiphase
boundary also causes a relaxation and a slight
bending of the atomic (0 0 1) planes close to the
antiphase boundary.
The interface stacking on the left side of the
antiphase boundary was determined to be: bulk±
SrO±TiO2 ±BaO±CuO±BaO±CuO2 ±Dy±CuO2 ±BaO±
bulk. For the stacking of the layers at the interface
on the right side of the antiphase boundary the
following sequence was found: bulk±SrO±TiO2 ±
BaO±CuO2 ±Dy±CuO2 ±BaO±CuO±BaO±bulk. Image simulations for both interface models were
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S. Bals et al. / Physica C 355 (2001) 225±230
Fig. 3. Di€erent types of possible interface arrangements: (a)
bulk±SrO±TiO2 ±BaO±CuO2 ±Dy±CuO2 ±BaO±CuO±BaO±bulk,
(b) bulk±SrO±TiO2 ±BaO±CuO±BaO±CuO2 ±Dy±CuO2 ±BaO±
bulk, (c) bulk±SrO±TiO2 ±Dy±CuO2 ±BaO±CuO±BaO±CuO2 ±
Dy±bulk.
performed and the contrast could be matched
with the experimental images, both for contrast
and for position, a focus value of 24 nm and a
thickness of 2.4 nm was used for the left part of the
image while a thickness of 3.2 nm was used for the
right part. These simulated images are shown in
Fig. 2 as insets.
These two interface con®gurations were observed throughout the complete ®lm. The third
possible interface stacking: bulk±SrO±TiO2 ±Dy±
CuO2 ±BaO±CuO±BaO±CuO2 ±Dy±bulk, was never
observed.
4. Discussion
Ultra-thin ®lms of DyBa2 Cu3 O7±x prepared by
PLD on TiO2 -terminating SrTiO3 apparently contain numerous antiphase boundaries. Such boundaries were already observed in REBa2 Cu3 O7±x thin
®lms and in these cases they were related to unit
cell steps on the substrate. The miscut angle of the
substrate used in this study was less than 0.2°,
implying that the minimum width of a substrate
terrace is 112 nm. Since HREM showed that the
average width of an antiphase domain is of the
order of 20 nm, unit cell steps at the substrate can
impossibly be the only reason for the formation of
antiphase boundaries. Antiphase boundaries here
are necessarily created during the growth of the
®lm and the consequence of a varying stacking
mode at the interface. Such antiphase boundaries
will in¯uence the superconducting properties, because 1 out of 2 superconducting CuO2 planes will
be disconnected.
The ®rst layer of the DyBa2 Cu3 O7±x thin ®lm
was determined to be a BaO layer. Since the substrate has a B site termination, the ®rst layer of the
®lm is expected to be an AO layer in order not to
disturb the perovskite stacking. The possibility of
a BO2 layer on a BO2 layer is hardly ever observed
[4,13]. An interface with BaO as a ®rst layer of the
®lm is apparently more stable than one with Dy as
®rst layer. The latter con®guration was never observed.
The instability of an interface of the type C
(Fig. 3) can be understood by taking into account
the co-ordination environment of the Ti4‡ cations. The SrO±TiO2 ±Dy±CuO2 stacking sequence
would imply an incomplete square-pyramid polyhedron for the Ti cation whereas it is favourably
disposed towards an octahedral co-ordination.
The absence of oxygen atoms in the Dy layer also
results in an increase of the lattice energy associated with the interface due to a repulsion between
the highly charged Ti4‡ and Dy3‡ cations. As a
semiquantitative measure of the lattice energy and,
hence, the stability of di€erent interfaces one can
use the Madelung constant a. It should be noted
that an absolute value of the electrostatic energy is
meaningless for HTSC compounds because of the
strong covalent bonding in the (CuO2 ) planes.
Nevertheless, it is possible to compare the a values for di€erent types of interface con®gurations
since the interaction between Ti4‡ , Dy3‡ , Ba2‡ and
oxygen can be considered as mostly ionic. The
Madelung constant was computed using the Ewald
method of convergent series with an algorithm
S. Bals et al. / Physica C 355 (2001) 225±230
described in Ref. [14]. The charges associated with
statistically occupied positions were chosen in accordance with the site occupancies. The unit cell
constructed for calculations consists of one unit
cell of DyBa2 Cu3 O7±x and one unit cell of SrTiO3
with a stacking allowing to keep the required layer
sequence across the interface and the periodicity in
the direction normal to the interface plane. The
computed a values are associated with these arti®cial structures, which contain an equal amount of
atomic layers of the same type and di€er only by
the layer sequences corresponding to the interfaces
A, B and C. We found that the a value is drastically lower for the interface C …a ˆ 52:05† than for
interfaces A …a ˆ 64:18† and B …a ˆ 63:85†. This
rough estimate suggests that the C interface is
unstable in comparison with the interfaces A and
B, which have similar lattice energies and, as a
consequence, (almost) equal formation probabilities.
Haage et al. [15] proposed a variable stacking
mode, with the possibility of partial overgrowth at
substrate steps, which would favour step-¯ow
growth. They studied a YBa2 Cu3 O7±x thin ®lm
deposited on a (1 0 6) SrTiO3 substrate and they
found that the terrace width of the ®lm is larger
than that of the substrate. In our study a substrate
with a much lower miscut angle was used and we
found far more antiphase boundaries than unit cell
steps at the substrate, which means that the terraces of the substrate are wider than the terraces of
the ®lm. It seems probable that two di€erent interface stacking modes can be formed because the
two possible interface sequences have comparable
energies and, consequently, antiphase boundaries
are formed. We support the existence of a variable
stacking sequence; however it is clear that this induces the formation of extra antiphase boundaries
when a substrate with a low miscut angle is used.
From our experiments we propose the following
growth sequence. Initially pyramidal shaped islands are formed. Because of the TiO2 terminated
substrate, the ®rst layer of each island is BaO. The
next layer however can be di€erent in neighbouring pyramidal islands; it can be either a CuO2 or a
CuO layer. When neighbouring islands coalesce,
antiphase boundaries can be created, all of the
same type with the same displacement vector.
229
5. Conclusion
We have studied the initial growth of ultra-thin
DyBa2 Cu3 O7±x ®lms on a TiO2 terminated SrTiO3
substrate. At the interface a variable atomic
stacking sequence was found. Two di€erent interface arrangements were determined: bulk±SrO±
TiO2 ±BaO±CuO±BaO±CuO2 ±Dy±CuO2 ±BaO±bulk
and bulk±SrO±TiO2 ±BaO±CuO2 ±Dy±CuO2 ±BaO±
CuO±BaO±bulk. This variable stacking sequence will
cause a structural and chemical mis®t and will lead to
the formation of antiphase boundaries, which will
in¯uence the superconducting properties of the ®lm.
The growth of DyBa2 Cu3 O7±x ®lm proceeds following a growth mode based on step-¯ow growth, but
with islands acting as terraces rather than steps at the
substrate. Due to an almost equal formation probability, both mentioned interface arrangements occur, leading to a dramatic increase of antiphase
boundaries.
Acknowledgements
Part of this work has been performed within the
framework of IUAP 4/10. The authors are grateful
to Dr. A. Abakumov of Moscow State University
for scienti®c discussions about the chemical aspects of the problem and to Dr. D.V. Fomitchev
for permission to use his program for calculating
the Madelung constant. S. Bals is grateful to the
Fund for Scienti®c Research-Flanders (FWO),
Belgium.
References
[1] H.W. Zandbergen, J.G. Wen, C. Traeholt, V. Svetchnikov,
J. Alloys Compounds 195 (1993) 85.
[2] B. Dam, B. Stauble-P
umpin, J. Mater. Sci. 9 (1998) 217.
[3] N. Missert, R. Hammond, J.E. Mooij, V. Matijasevic,
P. Rosenthal, T.H. Geballe, A. Kapitulnik, M.R. Beasley,
S.S. Laderman, C. Lu, E. Garwin, R. Barton, IEEE Trans.
Magn. 25 (1989) 2418.
[4] J.G. Wen, C. Traeholt, H.W. Zandbergen, Physica C 205
(1993) 35.
[5] J.N. Eckstein, I. Bozovic, Annu. Rev. Mater. Sci. 25 (1995)
679.
230
S. Bals et al. / Physica C 355 (2001) 225±230
[6] V.C. Matijasevic, B. Ilge, B. Stauble-Pumpin, G. Rietveld,
F. Tuinstra, J.E. Mooij, Phys. Rev. Lett. 76 (1996)
4765.
[7] B. Roas, L. Shultz, G. Endres, Appl. Phys. Lett. 53 (1988)
1557.
[8] T. Haage, J. Zegenhagen, H.-U. Habermeier, M. Cardona,
Phys. Rev. Lett. 80 (1998) 4225.
[9] G.J.H.M. Rijnders, G. Koster, D.H.A. Blank, H. Rogalla,
Appl. Phys. Lett. 70 (1997) 1888.
[10] R. Ramesh, A. Inam, D.M. Hwang, T.S. Ravi, T. Sands,
J. Mater. Res. 6 (1991) 2264.
[11] J. Zegenhagen, T. Siegrist, E. Fontes, L.E. Berman, J.R.
Patel, Solid State Commun. 93 (1995) 763.
[12] G. Koster, B.L. Kropman, G.J.H.M. Rijnders, D.H.A.
Blank, H. Rogalla, Appl. Phys. Lett. 73 (1998) 2920.
[13] G. Koster, Ph.D. Thesis, Arti®cially layered oxides by
pulsed laser deposition, University of Twente, ISBN
9036513367, 1999.
[14] S.G. Popov, V.A. Levitzkiy, Zh. Phys. Khim (Russ.) LV,
87 (1981) 87.
[15] T. Haage, J. Zegenhagen, J.Q. Li, H.-U. Habermeier, M.
Cardona, Phys. Rev. B 56 (1997) 8404.