PAPER www.rsc.org/materials | Journal of Materials Chemistry

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Synthesis, structure and properties of the hexagonal perovskite,h-BaTi1�xHoxO3�x/2

Yang Liu, Emma E. McCabe, Derek C. Sinclair and Anthony R. West

Received 23rd December 2008, Accepted 1st May 2009

First published as an Advance Article on the web 11th June 2009

DOI: 10.1039/b822785c

The crystal structure of hexagonal (h)-BaTi0.85Ho0.15O2.925 has been determined using neutron powder

diffraction data. The structure is derived from that of h-BaTiO3 and contains Ho3+, the largest cation

known to be accommodated by the B site in h-BaTiO3. Ti and Ho are disordered over the B1 octahedral

sites and the structure may be regarded as intermediate between those of h-BaTiO3 and

Ba3Sr(Nb,Ta)2O9. h-BaTi0.85Ho0.15O2.925 forms as a long-lived but metastable, intermediate phase

before transforming, slowly, to the thermodynamically stable cubic polymorph of the same

composition; its formation is an example of Ostwald’s rule of successive reactions. It is an electrical

insulator with relative permittivity of �50.

Introduction

The perovskite BaTiO3 exists in a number of polymorphs, all

related to the cubic ABO3 perovskite structure, also known as the

3C-type perovskite structure.1 This 3C-type structure consists of

close-packed AO3 layers stacked along [001] in a cubic sequence

(i.e. ABCABC) with B cations occupying octahedral sites to form

a lattice of corner-linked BO6 octahedra.2 The tetragonal poly-

morph of BaTiO3, in which the Ti4+ cations are displaced from

the centre of the octahedra towards an apex, has received much

attention as a result of the high permittivity at the tetragonal–

cubic phase transition (3max �10 000, TC �130 �C).3

Undoped BaTiO3 is stable in the 6H hexagonal polymorph at

temperatures above �1460 �C and adopts this structure until it

melts at�1620 �C.4 The structure of the 6H (here also referred to

as h) polymorph of BaTiO3-related materials is composed of 6

close-packed BaO3 layers stacked in the sequence [cch]2, where c

and h refer to cubic and hexagonal stacking, respectively. This

results in two different octahedral sites for the B cations: B(1)

occupies corner-linked octahedra and B(2) is located in face-

shared B(2)2O9 dimers.5 The hexagonal polymorph of undoped

BaTiO3 can be stabilised to room temperature by heating BaTiO3

in reducing conditions, resulting in partial reduction of Ti4+ to

Ti3+ and the creation of oxygen vacancies in O(1) sites in the

hexagonal layers;6 the general formula is BaTi4+1�xTi3+

xO3�x/2.

In addition, a number of B site dopant cations, of comparable

size to Ti4+, can stabilise the 6H polymorph to lower tempera-

tures. For aliovalent substitution of lower valence cations, charge

compensation is again achieved by creation of oxygen vacancies

in the O(1) sites.7–12 The permittivity of doped and undoped

h-BaTiO3 at room temperature is typically in the range 20 # 30 #

100, giving rise to possible microwave dielectric applications.13

Various mechanisms for stabilisation of this 6H polymorph

have been put forward yet it is still not well understood. Dickson

et al. proposed that transition metal cations with partially filled

d orbitals are necessary, with stabilisation driven by the

Department of Engineering Materials, University of Sheffield, MappinStreet, Sheffield, UK S1 3JD

This journal is ª The Royal Society of Chemistry 2009

formation of metal–metal bonds within the B(2)2O9 dimers.7 Ren

et al. proposed that cations such as Mn4+, comparable in size to

Ti4+, are essential.14 Langhammer et al. suggested that Jahn–

Teller active dopants (especially Ti3+, Mn3+ and Cu2+) are

required.15,16 However, stabilisation with p-block cations8,9 such

as Ga3+ also occurs and indicates the need for continued inves-

tigation of the stabilisation mechanism of this high temperature

phase.

Here, we report the structural characterisation of h-BaTiO3

stabilised kinetically by partial substitution of a much larger

cation, Ho3+, in place of Ti4+, together with its electrical prop-

erties. The phase—h-BaTi0.85Ho0.15O2.925—forms as a long-lived

metastable intermediate before formation of the thermodynam-

ically stable cubic (c)-polymorph of the same composition. The

c-polymorph is then stable at all temperatures and, once formed,

does not transform back to the h-polymorph. This work forms

part of a larger study into the phase equilibria, solid solution

mechanisms and electrical properties of Ho-doped BaTiO3, to be

reported elsewhere.17

Experimental procedure

Polycrystalline samples of h-BaTi0.85Ho0.15O2.925 were prepared

by solid state reaction of stoichiometric quantities of dried

BaCO3, TiO2 and Ho2O3. The reagents were intimately ground in

an agate mortar and pestle, heated at 1050 �C for 12 h and then at

1550 �C for 24 h in air in Pt crucibles, with intermittent grinding.

Pellets for electrical characterisation were prepared by heating

the reagents at 1050 �C for 12 h, then at 1450 �C for 12 h, pressed

into pellets using a uniaxial press followed by a cold isostatic

press (CIP) and sintered at 1550 �C for 12 h. This gave pellets

with density of�74% of the theoretical value. It was not possible

to increase the pellet density by firing at higher temperatures

because the structure then changed (slowly) to that of the

c-polymorph. Electrodes were fabricated on opposite pellet faces

using Au paste which was dried, decomposed and hardened by

gradually heating to 800 �C. Impedance spectroscopy (IS)

measurements were carried out using a Hewlett Packard 4192A

Impedance Analyser between room temperature and 500 �C in

J. Mater. Chem., 2009, 19, 5201–5206 | 5201

Table 2 Selected bond lengths and inter-ionic distances in �A

Ba(1)–O(1) 3 � 2.9118(1) Ti/Ho(1)–O(2) 6 � 2.0713(6)Ba(1)–O(1) 3 � 2.9116(1) Ti(2)–O(1) 3 � 2.050(1)Ba(1)–O(2) 6 � 2.9270(6) Ti(2)–O(2) 1 � 1.923 (1)Ba(2)–O(1) 3 � 2.855(1) Ti(2)–O(2) 2 � 1.923(1)Ba(2)–O(2) 2 � 2.9144(1)Ba(2)–O(2) 1 � 2.9143(1) Ti(2)–Ti(2) 2.818(2)Ba(2)–O(2) 2 � 2.9141(1) Ti/Ho(2)–Ba(2) 3.4383(4)Ba(2)–O(2) 1 � 2.9141(1) Ti/Ho(2)–Ba(2) 3.4384(4)Ba(2)–O(2) 3 � 3.0801(9) Ti/Ho(2)–Ba(2) 3.572(1)

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the frequency range 10 Hz to 1 MHz. X-Ray powder diffraction

(XRPD) data were collected at room temperature over a period

of 10 h using a Sto€e STADI P diffractometer operating in

transmission mode with a monochromated Cu Ka1 radiation

source and a position sensitive detector with a step size of 0.01�

2q. Time of flight (TOF) neutron powder diffraction (NPD) data

were collected at room temperature on GEM at ISIS, Oxford,

UK. Rietveld refinements used the GSAS suite of programs.18

Selected area electron diffraction (SAED) patterns were taken

using a Philips EM430 transmission electron microscope, accel-

erating voltage 300 kV.

Results and discussion

Structural characterisation of h-BaTi0.85Ho0.15O2.925

Structural refinement. Both the XRPD and NPD patterns were

indexed on a hexagonal unit cell, a z 5.81 �A and c z 14.31 �A

with space group P63/mmc. A refinement was carried out using

the structure reported by Akimoto et al.5 as a starting model. The

lattice parameters were first refined using the bank 6 data and

were then fixed as subsequent banks of data were added to the

refinement. Background (shifted Chebyshev function), histo-

gram scale factor and peak profiles were refined for each histo-

gram. Atomic coordinates were refined with constraints applied

to the temperature factors and positions of Ti and Ho on the 2a

site. In the initial model, Ti and Ho were distributed statistically

across both B sites but this model did not give a sensible

refinement. Subsequent results showed that the B(2) site was

occupied exclusively by Ti and this distribution was therefore

fixed. The fractional occupancies of both oxygen sites were

initially refined freely giving an overall oxygen content of

2.869(5) per formula unit, with vacancies located predominantly

on the O(1) site in the hexagonal layer. This is somewhat less than

the expected value of 2.925 per formula unit. However, when the

overall oxygen content was fixed to 2.925, there was a negligible

change in agreement between the calculated and observed

patterns. The oxygen content is related directly to the oxidation

states of the cations. The only cation capable of variation in

oxidation state is Ti. However, Ti3+, even in small concentration,

causes the electrical conductivity of titanates to rise dramatically.

Since the present phase is highly insulating (see later), the Ti must

be present as Ti4+ and hence the oxygen content is given directly

by the cation stoichiometry. In the final refinement, the overall

oxygen content was fixed at 2.925 and the occupancies of O(1)

and O(2) refined within this overall constraint. Final parameters

Table 1 Structural parameters from refined room temperature TOF NPD d

Atom Wyckoff site x y

Ba(1) 2b 0 0Ba(2) 4f 1/3 2/3Ti/Ho(1) 2a 0 0Ti/Ho(2) 4f 1/3 2/3O(1) 6h 0.51869(9) 0.0374(1)O(2) 12k 0.83130(7) 0.6626(1)

a a ¼ 5.8112(2) �A, c ¼ 14.2830(8) �A, c2 ¼ 14.89 (81 variables), Rwp ¼ 3.76%

5202 | J. Mater. Chem., 2009, 19, 5201–5206

are given in Table 1, selected bond lengths in Table 2, refinement

profiles in Fig. 1 and the structural model in Fig. 2.

h-BaTi0.85Ho0.15O2.925 adopts a B site ordered 6H structure in

which the face-shared B(2)O6 octahedra are occupied exclusively

by Ti, whilst the B(1)O6 octahedra between the c-stacked BaO3

layers accommodate both Ti and Ho. The O vacancies are

located in both the h-BaO3 layers and the c-BaO3 layers.

P63/mmc vs. P63/m. The ideal (hcc)2 structure has space group

P63/mmc but materials such as Ba3SrNb2O9 and Ba3SrTa2O9

adopt slightly distorted structures with P63/m symmetry. The

lower symmetry is manifested by tilting of the octahedra around

one of their three-fold axes. These two phases show complete

ordering of Sr and Nb/Ta over the two B sites: Sr occupies the

larger cubic B(1) site and Nb/Ta are located in the B(2)2O9 face-

shared octahedra. Due to the substantial difference in size of Sr2+

and Nb5+/Ta5+ (six-coordinate ionic radii 1.18, 0.64 and 0.64 �A,

respectively),19 stacking of the different-sized octahedra is facil-

itated by rotation of the octahedra.20

The structure of these phases shows similarities with

h-BaTi0.85Ho0.15O2.925: both contain relatively large, lower val-

ent cations substituted onto the B(1) sites. The c/a ratio of

2.4578(2) observed for h-BaTi0.85Ho0.15O2.925 is slightly larger

than the accepted maximum of 2.45 for structures of P63/mmc

symmetry, above which distortions resulting in structures of P63/

m symmetry usually occur. A high c/a ratio, 2.47, is also observed

for BaFe0.67Ti0.33O2.664 of P63/mmc symmetry,10 but the oxygen

vacancies may give rise to an expansion which might account for

the large c/a ratio.

Structural refinements for h-BaTi0.85Ho0.15O2.925 were also

carried out using models of P63/m symmetry but no improve-

ment in fit was observed (Rwp of 3.20% with 51 variables,

compared with Rwp of 3.18% for the equivalent model of P63/

mmc symmetry with 48 variables). In the absence of additional

reflections in XRPD, NPD and SAED patterns due to loss of the

ata for h-BaTi0.85Ho0.15O2.925, space group P63/mmca

z Fractional occupancy Uiso � 100/�A2

1/4 1 0.68(3)0.09870(7) 1 1.17(2)0 0.55/0.45 0.22(5)0.8487(1) 1/0 0.87(3)1/4 0.950(2) 1.24(2)0.08307(3) 0.987(1) 1.19(1)

, Rp ¼ 3.37%.

This journal is ª The Royal Society of Chemistry 2009

Fig. 1 Observed (+), calculated (�) and difference (bottom) profiles

from refinement using room temperature TOF NPD data for h-

BaTi0.85Ho0.15O2.925, bank 6 (a), bank 5 (b), bank 4 (c), bank 3 (d) and

bank 2 (e), overall c2 ¼ 14.89, Rwp ¼ 3.76%, Rp ¼ 3.37%.

This journal is ª The Royal Society of Chemistry 2009

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c glide plane from P63/mmc, there was no justification for

lowering the symmetry from P63/mmc. The slightly larger

c/a ratio for h-BaTi0.85Ho0.15O2.925 may reflect some local tilting

of the octahedra due to the presence of both large HoO6 and

small TiO6 octahedra in the same set of sites between the cubic

layers, but the long range structure is best described with

P63/mmc symmetry. The Ti/Ho(1) site must exhibit considerable

local deviation from the average size shown by the bond

distances in Table 2 as Ti4+ is much smaller than Ho3+. Never-

theless, we find no evidence of cation order or site distortions

associated with this set of sites.

B cation coordination environments. The presence of dopant

cations exclusively on the cubic B(1) sites is relatively unusual for

substituted h-BaTiO3. In most cases, either a statistical distri-

bution of B site cations is found (as observed for

BaTi1�yGayO3�y/28) or a preference for the B(2) site in the face-

shared octahedra (as found for BaTi0.667Ir0.333O37). The fact that

Ho3+ is not located in the B(2) sites is probably due to its large

ionic radius (0.901 �A);19 location in the B(1) site minimises

cation–cation repulsions which would otherwise be significant if

Ho was located in the face-shared B(2) sites. It also increases

substantially the mean radius of the B(1) site and this ensures

adequate separation of Ba(2) and B(2) ions on the cation sub-

lattice (see later). To the best of our knowledge, Ho3+ is the

largest cation known to substitute onto the B sites in h-BaTiO3,

the next largest being Pt2+ (ref. 7) and In3+ (ref. 21) (with six-

coordinate ionic radii of 0.8 �A).19

The variation in metal–oxygen bond lengths clearly demon-

strates the site preferences of the two different-sized B site

cations; bond lengths of 2.0713(6) �A for Ti/Ho(1)–O represent an

average between expected bond lengths of �1.976 �A for Ti–O

and �2.3 �A for Ho–O. In contrast, the Ti(2)–O bond lengths are

within the range expected for Ti–O bonds. These observations

are consistent with the B(2) site containing exclusively Ti4+

cations whilst Ho is located only on the cubic B(1) site. Inter-

estingly, Ti(2) forms 3 long bonds to O(1) sites and 3 short bonds

to O(2) sites. This difference in bond lengths highlights the

repulsion between the two Ti(2) ions located in the face-sharing

B(2)2O9 dimers, causing the Ti(2) ions to move apart towards the

cubic layers and away from the h-BaO(1)3 layer. This repulsion

may be further enhanced by the occurrence of some O(1)

vacancies which reduces the shielding between Ti(2) ions.

Oxygen vacancy distribution. All models used in Rietveld

refinements were slightly improved when the oxygen vacancies

were not confined to the O(1) site; refinements suggest vacancies

are located on both O(1) and O(2) sites. This is relatively unusual

as recent studies have demonstrated that even for high vacancy

concentrations, in most systems, vacancies are located exclu-

sively on the O(1) site;6 however, other exceptions are reported,

such as BaFeO2.79.22

A and B metal atom array in 6H-ABO3. In an elegant structural

analysis of 6H-Ba(Ti,Fe3+,Fe4+)O3-d Grey et al.11 demonstrated

the significance of considering the metal sublattice of the 6H

structure in an attempt to understand the factors that control its

stability, especially on incorporation of oxygen deficiency. They

noted that the largest variations in interatomic distances were

J. Mater. Chem., 2009, 19, 5201–5206 | 5203

Fig. 2 Unit cell of h-BaTi0.85Ho0.15O2.925 viewed down b axis, with unit

cell axes shown as solid lines in grey, Ba as large grey spheres, Ti/Ho(1)O6

octahedra in black and Ti(2)O6 octahedra in grey.

Fig. 3 Plot of Ba(2)–B(2)0 distances against average B cation ionic

radius, as shown by Grey et al.,11 with BaTi0.85Ho0.15O2.925 shown to be

similar to oxygen stoichiometric 6H materials.

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associated with pairs of B(2) ions and/or Ba(2) ions along [001].

Twinning of the BaB array occurs about the h-Ba(1)O(1)3 layers

and this produces two non-equivalent Ba and B atoms, Fig. 2.

In their solid solution, Grey et al.11 noted a short B(2)–Ba(2)

distance prior to any oxygen deficiency and that this distance was

essentially invariant at 3.40(1) �A on removal of oxygen from the

O(1) site across the solid solution range. This distance was noted

to be shorter than that expected for a BaTi alloy (�3.58 �A) and

therefore resistant to further contraction. The net result of

removal of O(1) atoms is displacement of both Ba(2) and B(2) in

the z-direction; however, this causes a reduction in Ba(2)–B(2)0

distance, Fig. 2. To avoid an unreasonably short Ba(2)–B(2)0

distance (dashed line Fig. 2) adequate separation of the

c-Ba(2)O(2)3 layers is required (see dotted lines). In h-BaTi0.85-

Ho0.15O2.925, this is aided by the presence of a large ion (Ho) on

the B(1) site as it resides in octahedral holes between adjacent

c-Ba(2)O(2)3 layers.

Our results for h-BaTi0.85Ho0.15O2.925 are in excellent agree-

ment with the description provided by Grey et al.11 Oxygen

deficiency is greater from the O(1) site, the Ti(2)–Ti(2) separation

is large (�2.818(2) �A), the Ba(2)–B(2) separation, 3.4384(4) �A is

>3.4 �A. The Ba(2)–B(2)0 separation (3.572(1) �A) is increased due

to the large Ho cation located exclusively on the B(1) site.

Grey et al.11 constructed a plot of Ba(2)–B(2)0 separation

against mean B site ionic radius for selected 6H phases and noted

a linear relationship for stoichiometric ABO3 phases, ranging

from BaCrO3 to Ba3YIr2O9. The Ba(2)–Ba(2)0 distance for

h-BaTi0.85Ho0.15O2.925 lies on this line, Fig. 3.

Thermodynamic considerations. h-BaTi0.85Ho0.15O2.925 is not

thermodynamically stable but forms as a kinetically stable

intermediate during reaction of starting materials:17 on

5204 | J. Mater. Chem., 2009, 19, 5201–5206

prolonged heat treatment, the structure transforms to that of

a cubic BaTiO3 solid solution. Phase diagram studies show that

the high temperature hexagonal phase of BaTiO3 is in fact

destabilised on doping with Ho, as shown by the increase in the

cubic to hexagonal phase transition temperature with increasing

Ho-content.17 As a consequence, the h-polymorph of BaTi0.85-

Ho0.15O2.925 could exist, hypothetically, as a thermodynamically

stable phase at very high temperatures, [1600 �C but this does

not happen in practice as samples melt prior to any possible

phase transition.

Given that h-BaTi0.85Ho0.15O2.925 is not stable thermody-

namically, it is relevant to consider the factors that influence its

formation and, once formed, its kinetic stability. Most phase

transitions are accompanied by an increase in enthalpy, DH, with

increasing temperature and therefore, from the free energy

relationship, DG ¼ DH � TDS, the phase transitions, and the

higher temperature polymorphs, must also have an increase in

entropy, DS. Indeed, high temperature polymorphs are ther-

modynamically stable at high temperatures only because the

magnitude of the �TDS term reduces their free energy to below

that of the low temperature polymorph(s). In the case of

BaTi0.85Ho0.15O2.925, the stable c-polymorph melts before

transforming to the h-polymorph17 and therefore, the magnitude

of T required for the h-polymorph to be thermodynamically

stable is inaccessible in the solid state. The h-polymorph, there-

fore, has both higher DH and higher DS than the c-polymorph at

all realistic temperatures.

Since h-BaTi0.85Ho0.15O2.925 has higher free energy than

c-BaTi0.85Ho0.15O2.925, it is thermodynamically metastable; we

may therefore regard it as a high energy reaction intermediate in

the pathway that leads finally to the thermodynamically stable,

cubic polymorph. It has higher entropy than that of the final

product; once formed, there is a high activation barrier for the

h- to c-transformation and therefore it has considerable kinetic

stability.

There are three aspects to be considered in the formation and

kinetic stability of a metastable phase such as h-BaTi0.85-

Ho0.15O2.925. First, why does the phase form at all? Clearly it

must have a high negative enthalpy of formation, with a high

lattice energy, in order to be kinetically stable at high tempera-

tures. It therefore forms in preference to the final, thermody-

namically stable product, because it has high entropy and in

particular, has an entropy intermediate between that of the

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disordered mixture of ions in the early stages of reaction and that

of the final product with the lowest entropy.

Second, what are the structural features of the metastable

phase that contribute to its high entropy? In the case of

h-BaTi0.85Ho0.15O2.925 the Ti/Ho(1) site must exhibit consider-

able local disorder which probably involves positional disorder

of either or both the cations and the surrounding oxygens, as

well as the occurrence of oxygen vacancies. A rule-of-thumb

guide to stable solid solution formation is that atom sizes

should be within 15% of each other for significant substitution

to occur. Using octahedral ionic radii: Ho3+, 0.901 �A and Ti4+,

0.605 �A,19 the Ho3+ cation is �50% larger than Ti4+; this size

difference is well beyond the limits of what may be regarded as

generally acceptable for Ho and Ti to be disordered over the

same set of crystallographic sites. We therefore identify the

Ti/Ho(1) site, and the surrounding oxygens, as the principal

source of high entropy in the structure. In particular, it seems

highly unlikely that any of the Ti/Ho–O(1) bonds have length

2.071 �A and that instead, this length represents an average of

significantly shorter bonds for Ti–O and much longer bonds

for Ho–O.

Third, why is the transformation to the thermodynamically

stable product slow? In this case, the h- to c-transformation

involves a change in the oxygen stacking sequence and in the

nature of the linkage of the (Ti,Ho)O6 octahedra. It is, therefore,

a major reconstructive transformation which involves the

breaking and reforming of strong Ti–O bonds.

The formation and kinetic stability of h-BaTi0.85Ho0.15O2.925

are an example of Ostwald’s rule of successive reactions, in which

reactions proceed through one or more metastable intermediates

before reaching the equilibrium product.23 In practice, this rule is

a guideline rather than a quantitative and verifiable rule.

Although often not stated explicitly, there are very many

examples in the reactions and formation of inorganic solids

where Ostwald’s rule may be cited. h-BaTi0.85Ho0.15O2.925 is yet

another example.

Fig. 4 (a) Impedance complex plane plot, (b) Z00/M00 combined spec-

troscopic plot and (c) capacitance C0 versus frequency, for h-BaTi0.85-

Ho0.15O2.925 ceramic at 399 �C.

Fig. 5 Arrhenius plot for h-BaTi0.85Ho0.15O2.925: total resistances (-)

and bulk resistances (,) extracted from the Z0 intercepts on Z* plots.

Electrical characterisation of h-BaTi1�xHoxO3�x/2

The electrical properties of h-BaTi0.85Ho0.15O2.925 ceramics

(pellets of �74% theoretical density) were measured using

impedance spectroscopy. Impedance complex plane plots,

Fig. 4a, and combined Z00/M00 spectroscopic plots, Fig. 4b,

indicate the presence of two components. The relative resistances

of the two responses in the complex plane plots did not change

with temperature, suggesting that they have the same activation

energy. The high frequency component has the smaller capaci-

tance value, 7 � 10�12 F cm�1; as it corresponds to the main peak

in the M00 plot, it is attributed to the bulk impedance of the

sample and also has the highest resistance (largest peak in the Z00

plot). The low frequency component has a capacitance of 3 �10�10 F cm�1 at 399 �C and is attributed to a thin layer compo-

nent which, from the similar activation energy of the resistance

values, appears to be associated with a constriction resistance.24

Values for the two resistances as a function of temperature are

presented as Arrhenius plots in Fig. 5.

The impedance data can also be represented as plots of

capacitance against frequency, as shown in Fig. 4c, and the

bulk capacitance value may be readily extracted from the

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frequency-independent high frequency plateau. The bulk

permittivity of h-BaTi0.85Ho0.15O2.925 is relatively high for

a non-ferroelectric compound (30 �50) and is consistent with

literature reports of h-BaTiO3 and other doped h-BaTiO3

phases.13

Conclusions

The crystal structure of h-BaTi0.85Ho0.15O2.925 has been deter-

mined. It is a hexagonal perovskite derived from the structure of

h-BaTiO3 and contains Ho3+, the largest cation known to enter

the h-BaTiO3 structure in the B sites. Smaller dopant cations

generally occupy partially the face-sharing octahedral sites,

whereas Ho is preferentially located on the other set of octahe-

dral sites that are corner-linked. The structure therefore shows

similarities to those of Ba3SrNb2O9 and Ba3SrTa2O9 in which the

larger Sr2+ cation occupies exclusively the corner-sharing octa-

hedral sites; h-BaTi0.85Ho0.15O2.925 may be regarded as an

intermediate between these structures and that of h-BaTiO3.

It is unusual for ions as different in size as Ti4+ and Ho3+ to

replace each other in solid solution formation. In this case, the

resulting phase is thermodynamically metastable and may be

regarded as an entropically stabilised high temperature poly-

morph which forms as an intermediate in the reaction pathway

that yields, eventually, the thermodynamically stable, cubic, Ho-

doped BaTiO3 polymorph. The source of the high entropy of the

h-structure is revealed by the Rietveld refinement results. The Ti/

HoO(1) octahedra contain two very different-sized cations,

whose average bond length to oxygen is either too long or too

short for the individual, cation–oxygen distances. Consequently,

there must be considerable local positional disorder of the atoms

involved, contributing to the high entropy of the structure.

The electrical properties of h-BaTi0.85Ho0.15O2.925 are similar

to those of other doped hexagonal perovskites; the permittivity is

high for a non-ferroelectric material with a value of �50 and the

phase is an electronic insulator with a conductivity of e.g. 2 mS

cm�1 at 600 K. The ceramics were not of high density, in spite of

the high sintering temperatures, as demonstrated by the imped-

ance data which indicated that the samples were electrically

inhomogeneous with bulk and constriction resistance compo-

nents with activation energies in the range 0.6–0.7 eV: constric-

tion resistances are a common feature of poorly sintered

ceramics.

5206 | J. Mater. Chem., 2009, 19, 5201–5206

Acknowledgements

We thank EPSRC for funding, Dr R. Smith for the collection of

NPD data, Dr H. Bagshaw for the collection of SAED patterns

and Prof. A. J. Bell (Leeds) for use of a CIP.

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