Structures and phase diagram for the system CaTiO3-La2/3TiO3

Article abstract of DOI:10.1016/j.jssc.2007.01.005

High-resolution neutron and synchrotron X-ray powder diffraction have been used to examine various compositions across the system (1-x)CaTiO₃-xLa_2/3TiO₃. The structures at room temperature were determined according to comp

Full text of DOI:10.1016/j.jssc.2007.01.005

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Journal of Solid State Chemistry 180 (2007) 1083–1092
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Structures and phase diagram for the system CaTiO –La TiO
3
3
2/3
a,
Ã
a,b
a
c
Zhaoming Zhang , Gregory R. Lumpkin , Christopher J. Howard , Kevin S. Knight ,
b,d
Karl R. Whittle , Keiichi Osaka
e
a
Australian Nuclear Science and Technology Organisation, Private Mail Bag 1, Menai, NSW 2234, Australia
b
Department of Earth Sciences, University of Cambridge, Downing Street, Cambridge CB2 3EQ, UK
c
ISIS Facility, Rutherford Appleton Laboratory, Chilton, Didcot OX11 0QX, UK
d
Department of Engineering Materials, University of Sheffield, Sir Robert Hadfield Building, Mappin Street, Sheffield S1 3JD, UK
e
Japan Synchrotron Radiation Research Institute/SPring-8, 1-1-1 Kouto, Sayo-cho, Sayo-gun, Hyogo 679-5198, Japan
Received 24 August 2006; received in revised form 21 December 2006; accepted 4 January 2007
Available online 14 January 2007
Abstract
High-resolution neutron and synchrotron X-ray powder diffraction have been used to examine various compositions across the system
(
1ꢀx)CaTiO
3
ꢀxLa2/3TiO . The structures at room temperature were determined according to composition: in Pbnm for 0pxp0.5,
3
Ibmm for 0.5oxo0.7, then I4/mcm for 0.7txo0.9, and finally in Cmmm for xX0.9. Although the four structures are the same as those
proposed previously in an X-ray diffraction study, the phase boundaries are somewhat different, in particular the Pbnm2Ibmm phase
boundary has been extended from xt0.4 to x40.5 in the current study based on our high-resolution neutron diffraction data. From in
situ measurements to identify structures above room temperature, an approximate composition–temperature phase diagram has been
¯
constructed, involving four temperature-induced phase transitions: Pbnm-I4/mcm, Ibmm-I4/mcm, I4/mcm-Pm3m and Cmmm-
P4/mmm.
r 2007 Elsevier Inc. All rights reserved.
Keywords: Perovskites; Crystal structures; Octahedral tilting; Cation/vacancy ordering; Phase diagram; Neutron powder diffraction; Synchrotron X-ray
diffraction
1
. Introduction
the perovskite A-sites, and the tilting of the TiO6
octahedra. At low La concentrations, hence low concen-
trations of vacancies (there is one vacancy for every two La
atoms in this system), the vacancies are thought to be
distributed at random, except perhaps for some association
Cation-deficient perovskites show high ionic conductiv-
ity indicating a considerable potential for use in solid oxide
fuel cells [1,2], and these materials exhibit unique dielectric
properties as well [3]. In spite of this, the crystal structures
of cation-deficient perovskites such as La2/3TiO3 and
Nd2/3TiO have until very recently been poorly understood
3
[4]. In consequence, pseudo-binary systems such as
SrTiO –La TiO and CaTiO –La TiO also have been
incompletely understood. The present study follows
elucidation of the crystal structure of La2/3TiO , as
stabilised by assorted dopants [5–7], and our own success-
ful study of the system SrTiO –La TiO [8].
The crystal structures in these cation-deficient perovs-
kites are determined by the disposition of vacancies across
3
+
of vacancies with La
ions. At high La/vacancy
concentrations, there is an energy advantage in the
accumulation of vacancies into alternate planes when
accompanied by the movement of the B-site (Ti) cations
toward and the anions away from these underpopulated A-
site planes [9]. There is ample experimental evidence, from
diffraction studies ([6] and references therein), that in the
La-rich compounds vacancies do indeed accumulate onto
alternate A-site planes, that is there is a layered ordering
involving the alternation of cation-rich and cation-poor
planes. At intermediate La concentrations, a degree of
layered ordering might be expected as in the SrTiO3–
3
2/3
3
3
2/3
3
3
3
2/3
3
Ã
La2/3TiO system [8]. Octahedral tilting in perovskites has
3
long been the subject of intensive study ([10] and references
Corresponding author. Fax: +61 2 9543 7179.
E-mail address: zhaoming.zhang@ansto.gov.au (Z. Zhang).
0022-4596/$ - see front matter r 2007 Elsevier Inc. All rights reserved.
doi:10.1016/j.jssc.2007.01.005

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Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1084
therein), and the different possible tilt systems are
described in terms of the component tilts around the x-,
y- and z-axes of the parent perovskite, making use of
Glazer’s notation [11,12]. Because of the corner connec-
diffraction data, we were able to prove that the correct
space group for the orthorhombic structure at x ¼ 0.4 and
0.5 is Pbnm, not Ibmm as previously reported. Vashook
and co-workers also carried out high-temperature studies
on a small number of compositions (x ¼ 0.4, 0.9, 0.96)
[18,19]. We have conducted studies on additional composi-
tions using both neutron and synchrotron X-ray diffraction
searching for phase transitions, so that a temperature–
composition phase diagram could be constructed.
tions of the BO octahedra the tilting of one octahedron
6
around any of these axes completely determines the tilting
in the plane perpendicular to that axis. The tilting of
octahedra in the next plane along the axis must be the same
in magnitude but can be either in the same (in-phase) or
opposite (out-of-phase) sense. Glazer’s symbol then is of
#
# #
the form a b c in which the literals refer in turn to tilts
around the parent x-, y- and z-axes and the superscript #
takes the values +, ꢀ, 0 to indicate in-phase, out-of-phase,
and no tilting, respectively. The repetition of a letter
indicates both tilt angles and lattice repeats, relating to the
corresponding axes, are equal. The crystal structures
corresponding to the different tilt systems with or without
layered cation/vacancy ordering have been obtained using
group theory and recorded elsewhere [10,13]. The experi-
mental determination of the crystal structures thus reduces
to the search for evidence for layered cation/vacancy
ordering and the nature of the tilts involved.
2. Experimental
Samples of Ca(
La2x/3TiO at x ¼ 0.1, 0.2, 0.3, 0.4,
1ꢀx)
3
0.5, 0.6, 0.7, 0.8 and 0.9 were produced by the standard
ceramic solid state technique. Stoichiometric quantities of
CaCO (Alfa-Aesar, 99.9%), La O (Aldrich, 99.5%) and
3
2
3
TiO (Alfa-Aesar, 99.9%) powders were mixed and ground
2
with ethanol in a ball mill. After drying, the mixtures were
heated in air at 1373 K for 1 day and then at 1523 K for 2
days, with cooling and regrinding after the initial heating
step. The crystallinity and phase purity of the products
were checked using laboratory powder X-ray diffraction
with Cu K radiation. The x ¼ 0 end member, CaTiO , was
The room temperature structure of the CaTiO3 end
member was established long ago [14]. It is orthorhombic
a
3
a high-purity powder sample purchased from Aldrich.
Time-of-flight powder neutron diffraction data were
recorded at room temperature for all samples using the
high-resolution powder diffractometer, HRPD, at the ISIS
neutron facility, Rutherford Appleton Laboratories, UK
[21]. The ground samples were loaded into thin-walled
11 mm diameter vanadium sample cans, which were then
suspended from the standard ISIS candlesticks. Diffraction
patterns from the samples were recorded over the time-of-
flight range 30–130 ms in both back-scattering and 901
detector banks, corresponding to d-spacings from 0.6 to
ꢀ
ꢀ +
in space group Pbnm, tilt system a a c . At 1498 K it
undergoes a discontinuous transition to a tetragonal
0
0 ꢀ
structure, space group I4/mcm, tilt system a a c , then at
634 K a continuous transition to the basic cubic structure,
1
space group Pm3m, tilt system a a a ([15] and references
0
0 0
¯
therein). As mentioned at the outset, the room temperature
structure for the other end member La2/3TiO has been
3
elucidated only very recently [5–7]. It is orthorhombic, in
ꢀ
0 0
space group Cmmm, and the tilt system is a b c . Upon
heating, there is a continuous transition as the tilt vanishes
ꢀ
4
˚
˚
(
but the layered ordering remains), and the structure at
elevated temperature is tetragonal, space group P4/mmm,
2.6 A (at a resolution Dd/dꢁ4 ꢂ 10 ) and from 0.9 to 3.7 A
ꢀ3
(Dd/dꢁ2 ꢂ 10 ), respectively. The patterns were normal-
ised to the incident beam spectrum as recorded in the
upstream monitor, and corrected for detector efficiency
according to prior calibration with a vanadium scan.
Patterns were recorded to a minimum total incident proton
beam of about 75 mA h at room temperature, correspond-
ing to ꢁ2.3 h of data collection, to allow good structure
determinations. High-temperature neutron diffraction pat-
0
0 0
tilt system a a c . The temperature of the orthorhombic to
tetragonal transition has been reported as 633 K in
La0.6Sr0.1TiO3 [6], 623 K in La0.68Ti0.95Al0.05O3 [16],
643 K in La0.63Ti0.92Nb0.08O3 [17], and 784 K in
La0.6Ca0.1TiO [18]. Vashook and co-workers [19,20] have
carried out mostly room temperature studies across the
3
(
and reported four structures according to composition: in
1ꢀx)CaTiO –xLa TiO system using X-ray diffraction,
3
2/3
3
terns were also obtained for Ca0.7La0.2TiO (x ¼ 0.3) up to
3
Pbnm (as for CaTiO ) for 0pxo0.4, then Ibmm for
1173 K, with the vanadium sample can mounted in an ISIS
designed furnace. The RAL furnace employs a cylindrical
vanadium element and operates under high vacuum
3
0
.4pxo0.7, I4/mcm for 0.7pxo0.8, and finally in Cmmm
for 0.8pxp0.96 (as for the end member La TiO ).
2
/3
3
ꢀ
4
However, this system covers a wide range of compositions
in pseudo-cubic symmetry (0.4pxp0.8), which means that
the correct structural identification relies on the presence of
superlattice reflections (arising from octahedral tilting in
this case). Since octahedral tilting involves the movement
of oxygen ions only, the corresponding superlattice
reflections would show more strongly in neutron diffrac-
tion patterns than X-ray ones. Therefore, we considered it
worthwhile to revisit this system using neutron diffraction
measurements. Based on our high-resolution neutron
(pressure o10 mbar). The thermometry is based on
type-K (chromel–alumel) thermocouples, the controlling
one positioned in contact with the sample can at about
20 mm above the beam centre. The sample temperature was
controlled to 70.5 K. The high-temperature patterns were
first collected in 100 K steps from 373 to 1173 K to a total
incident proton beam of about 75 mA h, then in 10 K steps
down to 1073 K to a total incident proton beam of about
35 mA h (finer temperature intervals but shorter counting
time in order to follow the phase transition in detail). In

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Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1085
order to allow for thermal equilibration of the sample, data
collection was delayed for a period of at least 5 min
depending on the size of the temperature step) after the
stability at the set point had been achieved.
High-temperature studies for other
Ca(
25
(
2
1
0
5
samples,
La2x/3TiO3 with x ¼ 0.4, 0.5, 0.8 and 0.9, were
x =0.9
.8
.7
1ꢀx)
0
carried out using the high-resolution X-ray powder
diffractometer installed on beam line BL02B2, at the
0
10
5
0
0
0
.6
.5
.4
SPring-8 synchrotron, Hyogo, Japan [22]. The wavelength
˚
was 0.7765 A as determined by calibration with a CeO
0.3
2
0
0
.2
.1
0
standard specimen. The finely ground samples were housed
in 0.2 mm quartz capillaries. All the diffraction patterns
were recorded to 751 (2y) on a single image plate. High
temperatures were achieved using a heated nitrogen gas
flow, and measured using a type-K thermocouple located
at the head of the nozzle. Patterns were recorded at
0
X
R
X
M
X
R
M
1.6
1.8
2.0
d-spacing (Å)
2.2
2.4
˚
Fig. 1. Extracts (1.6odo2.5 A) from the neutron diffraction patterns, as
(
sample) temperatures from 300 up to 980 K. The
temperature intervals were typically 50 K, but varied from
5 to 100 K depending on the perceived proximity to phase
recorded in the back-scattering detector bank, from samples of
3
Ca(1ꢀx)La2x/3TiO at room temperature. The groups of peaks are indexed
2
on the basic perovskite subcell—the brackets here are intended to refer to
a set of reflections that would be equivalent for the cubic perovskite.
Superlattice peaks at the M- and R-points arise from in-phase (+) and
transitions. For Ca0.3La0.47TiO3 (x ¼ 0.7), diffraction
patterns were obtained from 100 to 350 K with a
temperature interval of 50 K. Experimental details for the
low-temperature measurements are the same as those
described above, except that a glass capillary was used to
house the powder sample and the heated nitrogen gas flow
was replaced by a cooled one.
out-of-phase (ꢀ) tilting of the TiO
6
octahedra, respectively, while
superlattice peaks at the X-points arise either from the combined action
of in-phase and out-of-phase tilting, or cation/vacancy ordering on the
perovskite A-sites, depending on composition.
The neutron diffraction patterns were fitted using the
Rietveld method [23] as implemented in the GSAS
computer program [24,25]. Patterns from both the back-
scattering and the 901 detector banks were fitted simulta-
neously, the diffractometer constant for the 901 bank being
released to ensure that the lattice parameters were
determined by the higher-resolution back-scattering bank
in every case. The peak shapes were modelled as convolu-
tions of exponentials with a pseudo-Voigt in which two
peak width parameters were varied, and the background as
Chebyshev polynomials [24]. Displacement parameters
were refined along with internal co-ordinates, and oxygen
displacement parameters were taken to be anisotropic. The
isotropic displacement parameters for the two A-site
(x ¼ 0–0.9) at room temperature are shown in Fig. 1.
The peaks in the figure are identified by indices based on
the 1 ꢂ 1 ꢂ 1 parent perovskite subcell. Peaks with all
integral indices correspond to peaks from the ideal cubic
aristotype. The superlattice peaks are also indexed on the
basis of the 1 ꢂ 1 ꢂ 1 cell, and marked as R-point (all half-
integral indices), M-point (one integral, two half-integral
indices), and X-point (two integral, one half-integral
indices). Intensities at the R- and M-points arise from
out-of-phase (ꢀ) and in-phase (+) octahedral tilting,
respectively, and those at the X-point show contributions
either from the M- and R-point distortions acting in
concert [26], or from A-site cation/vacancy ordering. An
important first step in any data analysis is the inspection of
diffraction patterns, with a view to identifying the
structures obtained. The main perovskite peaks were
inspected for broadening or splitting, which would indicate
a tetragonal or orthorhombic distortion accompanying
cation ordering and/or octahedral tilting. In addition, a
search was made for superlattice peaks, at the R-, M- and
X-points.
2
+
3+
cations, Ca
and La , were constrained to be equal.
3
+
For Ca0.1La0.6TiO (x ¼ 0.9), the distribution of the La
/
3
2
+
Ca ions over the two crystallographically distinct A-sites
was allowed to vary, with the sum of the La and Ca
occupancies of these two sites constrained to 1.2 and 0.2,
respectively. Certain peaks seen to be broadened (M- or R-
point peaks—see next section) were distinguished using the
‘
[
stacking fault’ model within the GSAS computer program
24], then fitted with larger peak widths than the rest of the
Using the protocol described above, we identified the
structures as being orthorhombic in space group Pbnm, as
peaks.
for the CaTiO3 end member, for Ca
La2x/3TiO3 at
(1ꢀx)
x ¼ 0.1, 0.2, 0.3, 0.4 and 0.5. The tilt system is described in
ꢀ
ꢀ +
3
. Results and discussion
Glazer’s notation as a a c ; the presence of both in-phase
and out-of-phase tilting was indicated clearly by the
intensities at both M- and R-points in the diffraction
patterns. Weaker intensities were also observed at the X-
point, due to the combined action of in-phase and out-of-
phase tilting (not cation/vacancy ordering). Our results
3
.1. Room temperature studies
Extracts from the neutron diffraction patterns,
˚
La2x/3TiO3
1ꢀx)
1
.6odo2.5 A,
recorded
from
Ca(

ARTICLE IN PRESS
Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1086
differ from previous X-ray work [19,20], in which Ibmm
ꢀ
structures, respectively. Upon close inspection of the
˚
ꢀ 0
(tilt system a a c ) was assigned as the space group for
Ca( TiO3
doublet at dꢁ1.93 A (x ¼ 0.6 in Fig. 2(b)), indexed as
La2x/3TiO at x ¼ 0.4 and 0.5 (or Ca La
{200} on the parent subcell, it is clear that this peak has
p
1ꢀx)
3
x
2(1ꢀx)/3
at x ¼ 0.6 and 0.5), presumably because the M-point
reflections at these compositions were too weak to be
visible in the X-ray patterns. In contrast, the intensities at
higher intensity on the high-d side (i.e., a ob ¼ c ).
p
p
p
Therefore, we agree with the previously proposed structure
in Ibmm [19,20]. The similarity of diffraction patterns
across the boundary from Pbnm (apart from the vanishing
of the M-point reflections) also suggests that the transition
from that structure might be continuous, in which case the
structure in Ibmm would be the only possibility [10].
When x is increased from 0.6 to 0.7, the most visible
changes in the diffraction pattens are the decreased
intensities at the R-point reflections. However, the more
significant change occurs with the above mentioned
M-points in our neutron patterns for Ca(
La2x/3TiO at
3
1ꢀx)
x ¼ 0.4 and 0.5 were unmistakable, as illustrated in
Fig. 2(a). This confirms that neutron diffraction is
particularly valuable for monitoring octahedral tilting. In
addition, our high-resolution neutron diffraction data
revealed systematic broadening of the M- (and conse-
quently X-) point superlattice reflections for Ca(
La2x/3
1ꢀx)
TiO at x ¼ 0.2–0.5, the degree of broadening increasing
3
with x. Recalling that the M- point superlattice reflections
arise from in-phase tilting of the TiO6 octahedra, we
attribute this broadening to the occasional reversal of the
sense of this tilting, which could occur within an otherwise
unaltered framework and so not impact on the other peaks.
In other words, we believe this to be kind of a (domain) size
broadening, reflecting that the coherence of the in-phase
octahedral tilting is maintained over only a rather limited
range.
doublet {200} , which has collapsed into a singlet
p
(Fig. 2(b)). This means that the structure is pseudo-cubic
for Ca0.3La0.47TiO (x ¼ 0.7), and both Ibmm and I4/mcm
3
are now possible space groups. In other words, this
structure cannot be determined unambiguously based on
room temperature powder diffraction data alone. We
carried out refinements assuming both structural models,
and the goodness-of-fit is practically the same. However,
˚
the peak position of the R-point reflection at dꢁ2.33 A
For Ca0.4La0.4TiO3 (x ¼ 0.6), there was no longer
intensity at the M-point (Fig. 2(a)), this indicating a
different structure, without any in-phase octahedral tilting.
As there are no other qualitative changes in the diffraction
patterns, the structure should be either orthorhombic in
seems to be better described using the structural model in
I4/mcm. Therefore, we tentatively consider the room
temperature structure for Ca0.3La0.47TiO3 (x ¼ 0.7) as
tetragonal in I4/mcm and the phase boundary between
Ibmm and I4/mcm at just below this composition (xt0.7).
As shown in the next section, this assignment is consistent
with the phase diagram derived from high-temperature
measurements. It is worthwhile noting that another subtle
difference in the diffraction patterns between x ¼ 0.6 and
ꢀ
ꢀ 0
0 0 ꢀ
Ibmm (a a c ) or tetragonal in I4/mcm (a a c ). Based on
Glazer’s rigid octahedron model [11,12], the expected
relationships among the parent cell parameters are
a ob ¼ c and a ¼ b oc for Ibmm and I4/mcm
p
p
p
p
p
p
b
a
50
½{310}
x=0.6
x=0.5
1
1
1
4
{200}
x=0.9
x=0.8
x=0.7
2
0
8
6
4
2
0
40
30
20
x=0.4
x=0.3
x=0.6
x=0.5
x=0.2
x=0.1
x=0.4
x=0.3
10
x=0.0
0
2.38 2.40 2.42 2.44 2.46 2.48
1.90
1.92
1.94
1.96
d-spacing (Å)
d-spacing (Å)
˚
˚
Fig. 2. Extracts from the neutron diffraction patterns—as recorded in the back-scattering detector bank: (a) 2.38odo2.48 A, and (b) 1.90odo1.96 A,
showing the evolution of the 1 {310}
(M-point) and {200} reflections (indexed on the basic perovskite subcell) as a function of composition.
2
p
p

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Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1087
0
.7 is the broadened R-point peaks for x ¼ 0.7, suggesting
The pattern for Ca0.1La0.6TiO (x ¼ 0.9) is very similar
3
short-range order associated with the out-of-phase octahe-
dral tilting. Although no X-point reflections were observed
in the neutron diffraction patterns, they appeared as very
weak and broad features in the synchrotron X-ray patterns
at 300 K, indicating the onset of cation/vacancy ordering at
xE0.7. This also confirms that synchrotron powder XRD
is more sensitive to cation/vacancy ordering than neutrons.
The neutron diffraction pattern for Ca0.2La0.53TiO3
to that shown in our previous report for Sr La TiO [6].
0
.1
0.6
3
The structure was determined as orthorhombic in space
ꢀ
0 0
group Cmmm (tilt system a b c ), based on the splitting of
the main perovskite peaks [e.g., the {200} peak is now a
p
triplet as shown in Fig. 2(b)] and the presence of super-
lattice peaks at the X- and R-points (and very weak
intensities at the M-point due to the combination of X- and
R-point distortions). Pertinent features of this structure are
the alternation of (almost) fully and partly occupied layers
of La(Ca) cation, and out-of-phase tilting of the TiO6
octahedra around an axis perpendicular to the direction of
the cation ordering.
The neutron diffraction patterns were fitted using the
space groups identified above, and both lattice parameters
and atomic co-ordinates were determined using the
Rietveld method [23–25] as outlined in the previous
section. Details of the four distinct structures (represented
by four compositions at x ¼ 0.3, 0.6, 0.8 and 0.9), are
included in Table 1.
(
x ¼ 0.8) at room temperature is consistent with this
sample having a tetragonal structure in space group
I4/mcm, although the main perovskite peak still appeared
to be a single symmetric peak (Fig. 2(b)). This assignment
is also supported by the phase diagram obtained from high-
temperature measurements (see Fig. 7). Although we could
just make out the X-point reflections in the neutron
diffraction patterns (indicating the cation/vacancy ordering
has increased slightly with increasing La content—both in
degree and range), these peaks are still very weak and
broad and therefore not included in the refinements. In
contrast, Vashook and co-workers [20] reported that the
Fig. 3(a) shows suitably scaled cell parameters as a
function of composition (x) at room temperature (data for
x ¼ 0.7 were obtained assuming the tetragonal structure in
I4/mcm). The following phase boundaries are presented by
dashed vertical lines to indicate the change of space groups:
room temperature structure for Ca0.2La0.53TiO (x ¼ 0.8)
3
is orthorhombic in Cmmm, presumably including the weak
and broad X-point reflections in their structural refine-
ments.
Table 1
Lattice parameters, Wyckoff sites, atomic fractional coordinates, occupancies, and displacement parameters of the room temperature structures for
Ca0.7La0.2TiO
3
, Ca0.4La0.4TiO
3
3 3
, Ca0.2La0.53TiO and Ca0.1La0.6TiO
ꢀ
2
2
U (10 A )
˚
Atom
Site
x
y
z
Occupancy
Ca0.7La0.2TiO
3
˚
Pbnm, a ¼ 5.4285(1), b ¼ 5.4464(1), c ¼ 7.6851(1) A, Rwp ¼ 4.92%, R
p
¼ 4.65%
Ca/La
Ti
O1
4c
4a
4c
8d
0.0065(4)
0
0.5242(2)
0
0.25
0.7/0.2
1.35(2)
0.96(3)
1.39(9)
1.48(5)
0
0.25
1
1
1
ꢀ0.0623(3)
0.2158(2)
ꢀ0.0104(3)
0.2809(2)
O2
0.0311(2)
Ca0.4La0.4TiO
3
˚
Ibmm, a ¼ 5.4651(2), b ¼ 5.4637(1), c ¼ 7.7125(1) A, Rwp ¼ 6.32%, R
p
¼ 5.12%
Ca/La
Ti
O1
4e
4a
4e
8g
ꢀ0.0074(10)
0
0.5
0
0.25
0.4/0.4
1.43(3)
1.00(3)
2.08(16)
2.75(20)
0
0.25
1
1
1
ꢀ0.0454(5)
0.25
0
0.25
O2
0.0258(2)
Ca0.2La0.53TiO
3
˚
I4/mcm, a ¼ b ¼ 5.4709(1), c ¼ 7.7475(1) A, Rwp ¼ 7.28%, R
p
¼ 5.68%
Ca/La
Ti
O1
4b
4c
4a
8h
0
0.5
0
0.25
0
0.2/0.53
1.36(3)
1.32(4)
3.10(16)
1.90(7)
0
1
1
1
0
0.2227(2)
0
0.2773(2)
0.25
0
O2
Ca0.1La0.6TiO
3
˚
Cmmm, a ¼ 7.7489(1), b ¼ 7.7228(1), c ¼ 7.7733(1) A, Rwp ¼ 4.81%, R
p
¼ 4.08%
Ca/La 1
Ca/La 2
Ti
4g
4h
8n
4i
0.2526(4)
0.2578(9)
0
0
0
0.5
0.001 /0.928(3)
0.199 /0.272(3)
0.92(4)
1.16(11)
0.76(3)
0
0.2451(6)
0.2754(4)
0.2265(4)
0
0.2611(2)
0
0.5
1
1
1
1
1
1
O1
O2
0
1.24(15)
2.38(17)
1.16(18)
1.78(19)
1.84(10)
4j
0
O3
O4
4k
4l
0
0.2142(3)
0.2626(3)
0.2363(3)
0
0.25
0.5
0.25
O5
8m
The number in parentheses beside each entry indicates the estimated standard deviation referred to the last digit shown.

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Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1088
from Pbnm to Ibmm at 0.5oxo0.6, Ibmm to I4/mcm at
xE0.7, and I4/mcm to Cmmm at 0.8oxo0.9. Evidently
the lattice parameters of the basic perovskite subcell are
nearly equal in the composition range 0.4pxp0.8, similar
to data reported by Vashook et al. [20]. Although there
appears to be scatter in the reduced cell parameters, the
volume of the basic subcell (the volume per formula unit)
increases smoothly with composition, as shown in Fig.
From the atomic co-ordinates we can estimate the angles
of the out-of-phase (f) and in-phase (c) tilts of the TiO6
octahedra, using the formulae listed in Table 2. It should be
noted that in most structures (except I4/mcm) the out-of-
phase tilt angle (f) can be estimated from the positions of
either the apical or equatorial oxygen ions. The estimates
obtained are very similar, as they should be (in the absence
of shearing), and their average is presented in Fig. 4. As
shown, the two tilt angles decrease progressively with the
increase of La content. Note that the tilt angle for x ¼ 0.7
was essentially the same whether assuming the structure in
I4/mcm or Ibmm.
3
for Ca( La2x/3TiO was first reported by Kim et al. [3]
(b). The gradual increase of the basic cell volume with x
1ꢀx)
although no structural solutions were provided by these
3
(
authors), and it was attributed to the difference in ionic
3
+
2+
˚
˚
radii between La (1.36 A) and Ca (1.34 A), as well as
the creation of cation vacancies.
3.2. Variable temperature studies
a
The purpose of the variable temperature studies was to
3
3
3
3
3
3
.90
.88
.86
.84
.82
.80
c
/2
determine, for each composition, the transition tempera-
ture marking the onset (or disappearance upon heating) of
a particular octahedral tilting. Determination of the
transition temperatures by diffraction methods depends
on monitoring the corresponding superlattice reflections
and in some instances the splitting/broadening of the main
perovskite peaks, as summarised in Table 3. Among the
listed transitions, there are two first-order transitions
(Pbnm2I4/mcm and Ibmm2I4/mcm), and the rest are
likely to be continuous transitions [10,13]. It should be
noted that monitoring the main perovskite peaks for
splitting/broadening is a less-reliable way to determine the
c
/2
a/2
b/2
Pbnm
b/√2
/2
a/√2
c
Ibmm
I4/mcm Cmmm
a/√2
b
5
5
5
5
5
5
5
8.5
8.0
7.5
7.0
6.5
6.0
5.5
2
x
V
= -0.7593
+ 3.1471
x
+ 55.937
1
1
1
4
2
0
8
6
4
2
0
φ
ψ
Pbnm
Ibmm I4/mcm Cmmm
0.0
0.2
0.4
0.6
)
0.8
1.0
x
(La2/3TiO3
0.0
0.2
0.4
0.6
0.8
1.0
Fig. 3. Variation of (a) the reduced cell parameters, and (b) the basic
subcell volume, with composition (x) for Ca(1ꢀx)La2x/3TiO at room
x (La2/3TiO3)
3
temperature. The gap in the lines through the data points is intended to
reflect the necessarily discontinuous nature of the transition from Ibmm to
I4/mcm.
Fig. 4. Composition dependence of the octahedral tilting in Ca(1ꢀx)La2x/3
TiO at room temperature.
3
Table 2
Formulae for tilt angle calculations [27,28,6]
Structure, space group
Tilt axis
Out-of-phase tilt (f)
Tilt axis
[0 0 1]
In-phase tilt (c)
Orthorhombic Pbnm
[110]
[110]
p
tan fE4O2z(O2)
tan fEꢀ2O2x(O1)
tan fE4O2z(O2)
tan fEꢀ2O2x(O1)
tan f ¼ 1-4x(O2)
p
tan cE2[y(O2)-x(O2)]
Orthorhombic Ibmm
p
No in-phase tilting
Tetragonal I4/mcm
Orthorhombic Cmmm
[001]
[100]
p
No in-phase tilting
No in-phase tilting
p
tan fE2[z(O4)-z(O3)]
tan fE2[y(O1)-y(O2)]
Note that the lattice parameters have been removed from the formulae using aEbEc/O2 for Pbnm and Ibmm, and aEbEc for Cmmm.

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1089
Table 3
List of phase transitions encountered in the Ca(1ꢀx)La2x/3TiO
3
system, and associated changes in the diffraction patterns
Phase transition
Main changes in diffraction patterns
ꢀ
ꢀ
+
0
0
ꢀ
Pbnm (a a c )-I4/mcm (a a c )
First order
Disappearance of M-point reflections.
a
p
Splitting/broadening of the {h00} peaks into doublets with lower-intensity shoulders on the high-d side.
0
0
ꢀ
0 0 0
I4/mcm (a a c )-Pm 3¯ m (a a a )
Continuous
Disappearance of R-point reflections. No more splitting/broadening of the {h00}
peaks.
p
ꢀ
ꢀ
+
ꢀ
ꢀ
0
Pbnm (a a c ) - Ibmm (a a c )
Continuous
Disappearance of M-point reflections.
p
Splitting/broadening of the {h00} peaks into doublets with lower-intensity shoulders on the low-d side.
ꢀ
ꢀ
0
0
0
ꢀ
Ibmm (a a c )-I4/mcm (a a c )
First order
Intensities of the {h00} doublet reversed, and the splitting roughly doubled [29,30].
p
ꢀ
0 0
Cmmm (a b c )-P4/mmm
p
Disappearance of R-point reflections. The {h00} peaks change from triplets to doublets.
0 0 0
a a c )
(
Continuous
a
p 3
Note: the main {h00} peaks are not required by symmetry to be a singlet in Pbnm, however it appears to be an ‘‘accidental’’ singlet for many CaTiO
related materials.
a
b
3
2
2
1
1
0
5
0
5
0
5
0
2
1
1
0
.0
.5
.0
.5
½{321}
1
123 K
{200}
1
1
1
123 K
113 K
103 K
1113 K
1103 K
1
1
1
093 K
083 K
073 K
1093 K
1083 K
1073 K
2.04 2.06 2.08 2.10 2.12
1.92
1.94
1.96
d-spacing (Å)
d-spacing (Å)
Fig. 5. Extracts from the neutron diffraction patterns for Ca0.7La0.2TiO
1
3
(x ¼ 0.3)—as recorded in the back-scattering detector bank: (a)
and {200} reflections (indexed on the basic perovskite subcell) as
˚
˚
2
.025odo2.125 A, and (b) 1.91odo1.97 A, showing the evolution of the {321}
p
p
2
a function of temperature.
transition temperature for continuous phase transitions, as
the splitting/broadening would disappear at a temperature
below the actual transition due to limited instrument
resolution.
due to the coexistence of both Pbnm and I4/mcm phases),
and the main {200} peak became a doublet at the same
temperature. Therefore, we estimate the transition tem-
perature from the Pbnm orthorhombic structure to the
I4/mcm tetragonal one as roughly 1100 K for Ca0.7La0.2
p
Neutron (for x ¼ 0.3) or synchrotron (for x ¼ 0.4, 0.5,
0
.7, 0.8 and 0.9) diffraction patterns were recorded from
TiO (x ¼ 0.3). Similarly the transition temperature from
3
each sample at a number of temperatures, the temperature
range being selected according to the transition tempera-
ture anticipated. Figs. 5(a) and (b) show extracts from
Pbnm to I4/mcm is estimated, from X-ray synchrotron
diffraction measurements, to be around 850 K for Ca0.6
La0.27TiO (x ¼ 0.4).
3
1
neutron diffraction patterns for Ca0.7La0.2TiO (x ¼ 0.3),
Figs. 6(a) and (b) show the evolution of the {310} (M-
p
3
2
point) and {400} reflections in the temperature range of
1
2
corresponding to the
{321}p (M-point) and {200}p
p
reflections over the temperature range from 1073 to
123 K. As seen, the intensity of the M-point reflection
mostly disappeared at 1103 K (the residue intensity is likely
450–700 K for Ca0.5La0.33TiO (x ¼ 0.5), as determined by
3
1
synchrotron X-ray powder diffraction. While the intensity
1
of the {310} reflection mostly disappeared at 550 K, the
p
2

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Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1090
a
b
½{310}
{400}
3
3
2
2
.5
.0
.5
.0
7
00 K
15
10
5
7
6
00 K
50 K
6
6
50 K
00 K
6
00 K
5
50 K
550 K
500 K
5
00 K
450 K
450 K
0
18.0
18.2
18.4
18.6
47.0
47.5
48.0
2
θ (degrees)
2θ (degrees)
Fig. 6. Extracts from the synchrotron X-ray diffraction patterns for Ca0.5La0.33TiO
1
3
(x ¼ 0.5) from 450 to 700 K: (a) 17.9o2yo18.71, and (b)
4
6.8o2yo48.01, showing the evolution of the {310}
p p
(M-point) and {400} reflections (indexed on the basic perovskite subcell) as a function of
2
temperature.
shape and splitting of the main {400} peak did not show
p
obvious change until 100 K later, suggesting the presence of
two phase transitions: first from Pbnm to Ibmm at about
tetragonal structure in P4/mmm, the changes in the
synchrotron X-ray diffraction patterns are much more
obvious than for those samples in the pseudo-cubic range.
5
25 K then to I4/mcm at roughly 625 K. The intensity
The {h00} peaks, being triplets in Cmmm, became doublets
p
reversal of the {h00}p doublet and the (approximate)
doubling of the peak splitting, when the orthorhombic
Ibmm structure is transformed to the tetragonal I4/mcm
structure, are demonstrated clearly in Fig. 6(b).
at 700 K, indicative of a tetragonal structure. The intensity of
the R-point reflections vanished at the same temperature.
Therefore, the transition temperature from Cmmm to P4/
mmm is estimated as 675 K, which is about 100 K lower than
that reported by Vasylechko et al. (784 K) [18]. Due to the
large temperature interval used in their study (100 K), their
transition temperature was obtained by linear extrapolation
of the lattice parameters. Since the phase transition between
As described in the previous section, it is not possible to
distinguish unambiguously between the structures in Ibmm
and I4/mcm in the pseudo-cubic region at room temperature,
based on powder diffraction data alone. We have encoun-
tered the same problem in the variable temperature studies.
For Ca La TiO (x ¼ 0.7), synchrotron X-ray diffraction
1
Cmmm and P4/mmm may well be tricritical, linear extra-
polation could certainly overestimate the transition tempera-
ture. The other possible reason for the discrepancy is the
thermometry difference between different instruments.
Fig. 7 incorporates all the transition temperatures
derived from this study as well as those reported for
CaTiO (x ¼ 0) and Ca0.4La0.4TiO (x ¼ 0.6). The pub-
0.3 0.47
3
measurements were carried out from 100 to 350 K. The main
difference in diffraction patterns within that temperature
range is the slight narrowing of the main {h00} peaks with
p
increasing temperature (while the width of the {hhh} peaks
p
does not change). However, the {h00} peaks show as a single
p
3
3
(symmetric) peak even at 100 K, without any indication as to
lished transition temperatures for the CaTiO end member
3
the structure being in Ibmm or I4/mcm. Similarly, for
Ca La TiO (x ¼ 0.8), it is not possible to observe any
vary a great deal as reviewed recently by Ali and Yashima
[15]. We have used the values reported by these authors: at
0
.2 0.53
3
¯
splitting or asymmetric broadening of the main perovskite
peaks in the pseudo-cubic temperature range of 300–750 K
1498 K (Pbnm2I4/mcm) and 1634 K (I4/mcm2Pm3m).
For the Ca0.4La0.4TiO3 (x ¼ 0.6) sample, the published
(the sharpening of the {h00}p peaks with increasing
phase transition temperatures are around 450 K
¯
temperature was also observed). However, the determination
of transition temperature to the cubic structure was
straightforward, based on the disappearance of the R-point
reflections, at around 775 K for Ca La TiO (x ¼ 0.8).
(Ibmm2I4/mcm) and 1170 K (I4/mcm2Pm3m), respec-
tively [19]. While the Ibmm2I4/mcm phase transition is a
0.2 0.53
3
1
Although we did not recognise tricritical character in our initial study
of La0.6Sr0.1TiO [6], a similar transition in La NbO has been found to
be of tricritical form [31].
When Ca La TiO (x ¼ 0.9) undergoes its phase transi-
0.1 0.6
3
3
1/3
3
tion from the orthorhombic structure in Cmmm to the

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1091
4. Summary and conclusions
Compositions across the system (1ꢀx)CaTiO –
3
xLa2/3TiO3 have been studied using both neutron and
synchrotron X-ray powder diffraction techniques, at
various temperatures. The main conclusions have been
incorporated into the schematic phase diagram in Fig. 7.
Structures in this system are characterised by octahedral
tilting, as well as a layered cation/vacancy ordering on the
perovskite A-sites. In-phase (+) and out-of-phase (ꢀ)
octahedral tiltings are studied by monitoring the super-
lattice peaks, namely M-point (one integral and two half-
integral indices, when indexed on the basic perovskite
subcell), and R-point (all half-integral indices) reflections,
respectively. Diffraction patterns were recorded at each
composition as a function of temperature, in order to
estimate the temperatures of disappearance (upon heating)
or onset (upon cooling) of corresponding octahedral
tilting. The cation/vacancy ordering is evidenced by the
presence of peaks (X-point) with just one half-integral
index. These peaks are sharp, indicative of long-range
order for x ¼ 0.9, and a significant tetragonal distortion is
associated with this effect. The degree of cation/vacancy
ordering decreases with x (as evidence by much broader
and significantly less intense X-point reflections), with the
onset of short-range order at xE0.7.
Fig. 7. Phase diagram for Ca(1ꢀx)La3x/2TiO
3
based on transition
temperatures obtained from in situ studies of samples with different
compositions across the solid solution (solid dots are transition
temperatures determined in this study, while the solid and open squares
are previously published data [15,19]). The shading indicates the degree of
cation/vacancy ordering: random to xE0.7, then short range, and finally
long range for xX0.9.
(
to Pm3m is expected to be tricritical. Therefore, the
discontinuous) first-order one, the transition from I4/mcm
2
Based on high-temperature studies (Fig. 7), the room
temperature structures are orthorhombic in Pbnm for
0pxp0.5 and in Ibmm for 0.5oxo0.7, then tetragonal in
I4/mcm for 0.7txo0.9, and finally orthorhombic in
Cmmm for xX0.9. They are consistent with those
determined from room temperature neutron diffraction
patterns. It should be mentioned, however, that it is
difficult to determine the structures for x ¼ 0.7–0.8 based
on diffraction data alone, as they were found to be
metrically cubic to the resolution of our work. The
assignment of space groups for these two samples
(especially x ¼ 0.7) at RT is therefore only tentative, as it
relied upon the composition temperature phase diagram to
a large extent. Based on the refinement results from the
neutron diffraction patterns at room temperature, the
perovskite subcell volume increases, while both the in-
phase and out-of-phase octahedral tilt angles were found to
decrease, with increasing x. Both these trends are
consistent with the notion that the effective volume of
¯
¯
I4/mcm2Pm3m transition temperature of 1170 K would
be overestimated by Vashook et al. [19] using linear
extrapolation of the lattice parameters. In order to
compare their transition temperatures with those derived
from this study, we have shifted this transition temperature
for Ca0.4La0.4TiO by the same amount as the difference in
3
the transition temperature for Ca0.1La0.6TiO (see previous
3
paragraph).
The data points marking the observed transition
¯
temperatures between I4/mcm (Cmmm) and Pm3m
(
P4/mmm), corresponding to the transition from one out-
of-phase octahedral tilt to no tilt, can be fitted reasonably
well by a second-order polynomial function. The appar-
ently smooth variation of this transition temperature with
composition suggests that this transition, involving the
removal of tilt, is little affected by the degree of cation/
vacancy ordering. The phase boundary from Pbnm to
I4/mcm (or Ibmm), corresponding to the disappearance of
in-phase octahedral tilt, is also fitted by such a function.
Since we only have two experimental data points along the
Ibmm2I4/mcm phase boundary, linear extrapolation is
used as a (very) rough estimation. The degree of cation/
vacancy ordering, first random at low x, then short-range
starting at xE0.7, and finally long-range at xX0.9, is
indicated by shading in Fig. 7.
3
+
two La ions plus a vacancy is greater than the volume of
2+
the three Ca ions they replace.
Acknowledgments
We acknowledge the travel support to ISIS and Spring-8
for ZZ and CJH, provided by the Commonwealth of
Australia under the Major National Research Facilities
Program and the Australian Synchrotron Research Pro-
gram, respectively. The neutron facilities at ISIS are
operated by the Council for the Central Laboratory of
the Research Councils (CCLRC), with a contribution from
2
3
As, for example, the analogous transition in Ca0.3Sr0.7TiO [29].

ARTICLE IN PRESS
Z. Zhang et al. / Journal of Solid State Chemistry 180 (2007) 1083–1092
1092
the Australian Research Council (ARC). The synchrotron
facilities at SPring-8 are supported by the Japan Synchro-
tron Radiation Research Institute (Proposal No.
[16] M. Yashima, R. Ali, M. Tanaka, T. Mori, Chem. Phys. Lett. 363
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(
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[
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