Order—Disorder transformation and its effect on the properties of (Lanthanide)2Zr1.5Hf0.5O7 functional nanoceramics

Article abstract of DOI:10.1016/j.materresbull.2019.03.010

Order disorder transformations and its effect on the structural, optical and ionic transport properties of the cubic pyrochlore – defect fluorite structures are the prime focus of this study. Nanoceramics of Ln₂Zr_1.5Hf_0.5O₇ (Ln = Pr, Nd, Sm, Gd, Dy and Er) are prepared through the modified combustion technique. X – ray diffraction analysis and electron microscopic studies show the cubic nano crystallite nature of the materials. Vibrational studies reveal the pyrochlore–fluorite phase transformation in the series. Highly dense bulk ceramics are prepared from the nano particles at relatively low temperatures. The gradual replacement of Ln³⁺ ions in Ln₂Zr_1.5Hf_0.5O₇ changes the pyrochlores (Ln = Pr-Sm) to defect fluorite (Ln = Er) through the weakly ordered defect fluorites (Ln = Gd and Dy). Impedance spectroscopic analysis carried on the bulk ceramics shows the influence of phase transformation on the oxide ion conduction of bulk materials. The peak oxide ion conduction is observed for the Ln = Dy ceramic, which is at the pyrochlore fluorite phase boundary.

Full text of DOI:10.1016/j.materresbull.2019.03.010

Accepted Manuscript
Title: Order – disorder transformation and its effect on the
properties of (Lanthanide)2Zr1.5Hf0.5O7 functional
nanoceramics
Authors: Rejith R.S, Jijimon K. Thomas, Sam Solomon
PII:
S0025-5408(18)32389-4
DOI:
Reference:
https://doi.org/10.1016/j.materresbull.2019.03.010
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28 July 2018
4 March 2019
8 March 2019
Please cite this article as: Rejith RS, Thomas JK, Solomon S,
Order disorder transformation and its effect on the properties of
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(2019), https://doi.org/10.1016/j.materresbull.2019.03.010
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Order – disorder transformation and its effect on the properties of
(Lanthanide)2Zr1.5Hf0.5O7 functional nanoceramics
Rejith R.S, Jijimon K. Thomas and Sam Solomon*
Department of Physics, Mar Ivanios College, Thiruvananthapuram 695015
*Corresponding author: samdmrl@yahoo.com
1

Graphical abstract
Highlights
 Nanoceramics of Ln2Zr1.5Hf0.5O7 (Ln = Pr, Nd, Sm, Gd, Dy and Er) are prepared
through the modified combustion method.
 Structure gradually changes from the pyrochlore to the defect fluorite as the Ln ions
changes from Pr to Er.
 Hafnium substitution improves the densification of the bulk ceramics.
 Maximum ionic conductivity is observed at Pyrochlore – fluorite phase boundary.
ABSTRACT
Order disorder transformations and its effect on the structural, optical and ionic
transport properties of the cubic pyrochlore – defect fluorite structures are the prime focus of
this study. Nanoceramics of Ln2Zr1.5Hf0.5O7 (Ln = Pr, Nd, Sm, Gd, Dy and Er) are prepared
2

through the modified combustion technique. X – ray diffraction analysis and electron
microscopic studies show the cubic nano crystallite nature of the materials. Vibrational
studies reveal the pyrochlore–fluorite phase transformation in the series. Highly dense bulk
ceramics are prepared from the nano particles at relatively low temperatures. The gradual
replacement of Ln3+ ions in Ln2Zr1.5Hf0.5O7 changes the pyrochlores (Ln= Pr-Sm) to defect
fluorite (Ln=Er) through the weakly ordered defect fluorites (Ln=Gd and Dy). Impedance
spectroscopic analysis carried on the bulk ceramics shows the influence of phase
transformation on the oxide ion conduction of bulk materials. The peak oxide ion conduction
is observed for the Ln=Dy ceramic, which is at the pyrochlore fluorite phase boundary.
Key words: A. ceramics, B. Optical properties, C. Impedance spectroscopy, D.ionic
conductivity
1. INRODUCTION
The recent development in nano technology attracted the researchers to develop novel
materials with unique properties. In the field of ceramic processing, nanocrystalline ceramics
are expected to exhibit improved optical, magnetic, mechanical and electrical characteristics
due to their larger surface to volume ratio as compared to the conventional ceramic materials
[1]. Tailoring of these multifunctional properties in technologically important structures will
lead to multifunctional materials with improved characteristics. The A2B2O7 pyrochlore
structure was considered as a potential material in many fields due to the wide variety of
chemical substitution at the A (A2+ or A3+) and B (B5+ or B4+) sites [2]. Hence these are
widely used as catalysts, host for fluorescence centers, transparent ceramics, solid
electrolytes, nuclear waste encapsulates and thermal barrier coatings [3–7]. The unique
3

electrical properties make them potential candidates for solid oxide fuel cells (SOFCs)
applications.
Among the A2B2O7 Pyrochlores, Ln2Zr2O7 (Ln= lanthanides) were widely studied due
the variety of lanthanide elements at the Ln sites and the intrinsic oxide ion vacancies in the
structure. Depending on the cationic radius ratio and processing methods the Ln2Zr2O7
pyrochlores show different structural disorders [8]. Pyrochlore structure can be considered as
an ordered defect fluorite structure. In the Ln2Zr2O7 pyrochlore structure, the rare earth ion
Ln3+ is located at the 16c site, Zr4+ ion at the 16d site, and O2- oxide ions at 8a and 48f sites,
respectively. The 8b site is vacant in the completely ordered pyrochlore structure [9]. The
disorder in the structure increases as the oxide ions randomly occupies the 8b site. The
ordering process in the 8b site depends on the cationic radius ratio in the Ln2Zr2O7 pyrochlore
structure. By considering the 8 fold coordination to Ln3+ cations and 6 fold coordination to
Zr4+ cations, Ln2Zr2O7 compounds with cationic radius ratio (rA3+/rB4+) less than 1.46 will
form a disordered fluorite structure, between 1.46 and 1.78 will adopt an ordered pyrochlore
structure and greater than 1.78 it will have a layered perovskite structure [10]. That is, the
substitution of a larger cation in the A site or smaller cation in the B site on a disordered
fluorite structure will increase the oxide ion ordering, which in turn converts the defect
fluorite to pyrochlore structure. In this study the A site cation is successively replaced by the
Ln3+ ions and 25% of the B site cation is replaced by Hf4+ ions. Due to the fractional
substitution of Hf4+, only a slight change was expected in the structural parameters of the
well-known Ln2Zr2O7 pyrochlore structure, since the ionic radius of Hf4+ (0.71 Å) and Zr4+
(0.72 Å) have similar values in the 6 coordination [11]. However the slight changes in the
structural properties of the pyrochlore can create large amount of oxide ion vacancies in the
structure, which in turn affects the optical and ionic transport properties of the materials. It
was also reasonable to expect a different optical and ionic transport properties to the hafnium
4

substituted materials, since the electronic structure and density of Hf4+ ions are different from
Zr4+ ions.
The order - disorder transition in the pyrochlore structure is closely related to the ionic
transport properties of materials. H. Yamamura et al. studied the ionic transport properties of
Ln2Zr2O7 and found that the highest value of conductivity occurs at the vicinity of pyrochlore
– fluorite boundary [12]. Many researchers studied the order - disorder transitions by varying
A site and B site cations, a steady increase in conductivity was observed at the pyrochlore -
fluorite boundary [13–15]. The ionic transport properties of the dense ceramics also depend
on the bulk density and final heat treatment temperature, which in turn depend on the
preparation methods [16]. That is, there is an indirect dependency between the preparation
methods and the ionic transport properties of the dense ceramics. Preparation of dense
ceramic from the nanomaterials offers high density at low temperature as compared with the
conventional solid state method [17]. Therefore an auto ignited combustion method is used
for the preparation of nanoceramics in this study and then this powder is used for the
preparation of the final dense ceramics.
2. EXPERIMENTAL
In the combustion technique used for the preparation of Ln2Zr1.5Hf0.5O7 (Ln = Pr, Nd,
Sm, Gd, Dy and Er), the rare earth ions are prepared by dissolving their oxides (Pr6O11,
Nd2O3, Sm2O3, Gd2O3, Dy2O3 and Er2O3) in concentrated HNO3 whereas the Zr and Hf ions
are prepared by dissolving ZrOCl2.8H2O and HfCl4, respectively in double distilled water.
Then the stoichiometric amount of metal ions containing solution is mixed with citric acid
solution, keeping the citric acid to cation ratio unity. The precursor complex is prepared by
adding concentrated HNO3 (oxidant) and ammonia solution (fuel), keeping the pH ~7. This
5

O
solution is stirred well for uniform mixing and then heated in a hot plate at ̴ 250 C. The
solution boils and undergoes dehydration followed by decomposition leading to a smooth
deflation producing foam. On persistent heating the foam gets auto-ignited, giving a
voluminous fluffy powder due to self-propagating combustion. The obtained powder is
heated in the oxygen atmosphere at 850 OC for 30 minutes to remove the organic impurities.
The structure of the prepared nanomaterials are analysed using X-ray diffractometer
(D8 advance, Bruker, Germany) with CuKα radiation in the 2 range of 20–70 in steps of
0.02. Additional information regarding the phase purity of the powder samples are examined
using FTIR and confocal Raman studies. The infrared spectrum of the samples are recorded
in the range 370 – 1100 cm-1 on a FTIR Spectrometer (Spectrum2, Perkin-Elmer, Singapore)
by using the KBr method. Raman spectra of the powder samples are recorded between 200
and 1000 cm-1 using a confocal Raman spectrograph (alpha300RA, WITec GmbH, Ulm,
Germany) with a 532 nm DPSS laser. Particulate properties of the prepared nanoparticles are
imaged using high resolution transmission electron microscopy (Joel/JEM 2100, USA)
operating at 200 kV. The diffuse reflectance spectra in the range of 200 to 700 nm of the
samples are recorded by using a UV-Vis spectrometer (Lamda 35, PerkinElmer, Singapore)
with an integrated sphere accessory (RSA-PE-20, Labsphere, USA). In this method
Spectralon is used as the standard reference material for the samples. The photoluminescent
emission of the prepared nanopowders are recorded using a spectrofluorometer (FP-846,
Jasco, Japan). The phase pure powder samples are then ground in an agate mortar with
polyvinyl alcohol as the binding medium. It is then uniaxially compacted into pellets of 2 mm
thickness and 10 mm diameter at 20 MPa using a hydraulic press. The disc shaped pellets are
sintered in a high temperature resistive heating furnace with molybdenum heating elements
(TE-4050, Therelek, India). The relative density of the sintered pellets are calculated by using
Archimedes method. The surface morphology of the thermally etched pellet samples is
6

examined by Field Emission Scanning Electron Microscope (Nova NanoSEM 450).
Dielectric and impedance analysis are carried in an LCR meter (LCR HiTester 3532-50,
Hioki, Japan) in the frequency range 100 Hz – 5 MHz at different temperatures from 500 to
850 OC. Before taking the measurements, disc shaped capacitors are made by applying silver
paste on both sides of the pellets and fired up to 850 OC.
3. RESULTS AND DISCUSSION
The X-ray diffraction (XRD) patterns of the prepared Ln2Zr1.5Hf0.5O7 (Ln = Pr, Nd,
Sm, Gd, Dy and Er) nano ceramics are shown in Fig. 1. The XRD pattern of the Ln = Pr
nanoceramic shows the pyrochlore structure with space group Fd3m whereas all the other
prepared nano materials show the defect fluorite structure with space group Fm3m. ICDD
files 19-1021 and 78-1299 are respectively used for indexing the pyrochlore and defect
fluorite structures. The two low intensity super structure pyrochlore reflections (331) and
(511) nearly at the 2 values 37 and 45 degree, respectively are visible only in the XRD
pattern of Ln= Pr. The absence of these peaks in the other diffractograms may be due to the
broadening of the XRD peaks, which will results in the merging of low intensity peaks in the
background data. It is known that one of the major causes of peak broadening in the
diffraction pattern is the smaller size of the crystallites. Therefore the well-known Scherer’s
formula D = K/cos can used to determine the size of the nanocrystallites [18], where D is
the crystallite size, is the wavelength, is the full width half maximum (FWHM) of the
peaks and is the angle of diffraction. The calculated crystallite size values for the prepared
nanoceramics are shown in Table 1. It can be seen that the crystallite size decreases in the
Ln2Zr1.5Hf0.5O7 series from Pr to Gd and then increases from Gd to Er. Also it is identifiable
from the table that the lattice constant gradually decreases in the series as the lanthanides
7

changes from Pr to Er, which is due to the lanthanide contraction of Ln3+ ions. Here the
lattice parameters of fluorite structures are shown as doubled values for comparing it with the
pyrochlore structure.
Fig. 1.XRD patterns of Ln2Zr1.5Hf0.5O7 nanoceramics
8

Table 1. Structural parameters and crystallite size of Ln2Zr1.5Hf0.5O7 nanoceramics
Ln
ceramics
Ionic radius
(Å)
Crystallite size
(nm)
Lattice constant
(Å)
Pr
Nd
Sm
Gd
Dy
Er
1.126
1.109
1.079
1.053
1.027
1.004
8.93
6.45
4.96
4.80
5.73
6.09
10.70
10.69
10.64
10.58
10.48
10.42
X- ray diffraction sometimes fails to show the exact pyrochlore – fluorite transitions
in nanoceramics, since these structures are differed only by oxygen occupancy at the 8b site.
The intensity of the X-ray reflections from the oxygen atoms is very weak as compared with
the other metal atoms in the pyrochlore structure due to its relatively smaller size. Therefore
vibrational studies are employed in the prepared nano materials, the metal oxygen vibrations
are sensitive to the input frequencies of infrared and Raman studies. Factor group analysis
predicts 13 active vibrational modes in pyrochlore structure as [19]
Γ = A1g + Eg + 4F2g + 8F1u
For an ordered pyrochlore structure, the eight F1u modes are IR active (seven optical
and one acoustical) and the remaining six vibrational modes are Raman active. In the mid IR
region the absorption spectrum of pyrochlore structure is characterised by two absorption
bands whereas the fluorite structure is characterised by only one absorption band [17]. Fig.
2(a) shows the Fourier Transform Infrared (FTIR) spectra of the prepared materials between
370 and 1100 cm-1. The observed bands and the corresponding band assignments are
tabulated in Table 2. The FTIR spectra of the Ln2Zr1.5Hf0.5O7 with Ln = Pr – Sm shows two
9

distinctly visible absorption bands, Ln = Er shows only one absorption band and Ln = Gd and
Dy shows two faint broad absorption bands. That is Ln = Pr – Sm are purely pyrochlore
structured materials, Ln = Er is formed as a defect fluorite material whereas Ln = Gd and Dy
are defect fluorite structure with ordered pyrochlore domains. It is also clear from the table
(1)
that the F1u mode associated with the Zr/Hf- O bond gradually shift towards the lower
wavenumber region with the replacement of Ln ions from Pr to Er. However the F1u(2) mode
associated with the Ln - O' bond gradually shifted towards the lower wavenumber from Ln =
Pr to Ln = Sm and then shifted towards the higher wavenumber region from Ln = Gd to Ln =
Er. A broad shoulder peak around 600 cm-1 in the defect fluorite structures is an indication of
the presence of weakly ordered pyrochlore domains [16]. The broadening of peak around 600
cm-1 in Ln = Dy is higher than Ln = Gd, which indicates that the former one is less ordered
than the later.
Fig. 2 (a) FTIR spectra and (b) Raman spectra of the Ln2Zr1.5Hf0.5O7 ceramics
10

Table 2 Band assignments of FTIR and Raman spectra
FTIR spectrum
Raman spectrum
Eg
Sample
name
(1)
(2)
(2)
(3)
F1u
F1u
F2g (1) + F2g
F2g
Pr
Nd
Sm
Gd
Dy
Er
530 s
517 s
407 s
383 s
381 s
395 s
407 s
502 wbr
515 wbr
520 wvbr
386 mbr
375 mbr
392 mbr
398 wbr
294 s
310 s
505 s
490 s
330 sbr
335 vbr
481 sbr
420 sbr
Spectral intensity indications: s - strong, m – medium, w – weak and br - broad
The Raman spectra of Ln2Zr1.5Hf0.5O7 (Ln = Nd, Sm, Gd and Dy) are shown in Fig.
2(b). Here Ln = Pr and Er are avoided from the Raman spectral analysis due to the
fluorescence effect. The observed absorption bands and corresponding band assignments are
shown in Table 2. From the figure and table it is clear that Ln = Nd and Sm shows three
distinctly visible Raman active bands whereas Ln = Gd and Dy shows a similar spectra with
broad Raman bands. The peak broadening in the Raman spectra is not only influenced by the
nano size of the materials but also by the disorder in the structure. The translational
periodicity in the pyrochlore lattice are be broken down due to the presence of disorder in the
form of vacancy or defects in the crystal. This disordering allows more phonons to contribute
to the optical spectra. Therefore phonons from all parts of the Brillouin zones starts to
contributing to the vibrational spectra, which will creates broad, continuously spread weak-
intense Raman bands [20]. Thus six Raman active bands in a pyrochlore structure reduces to
a single broad Raman band around 350 cm-1 in the disordered defect fluorite structure [8].
However some defect fluorite structures usually shows weak intense peaks on their Raman
spectra due to the presence of short range ordered pyrochlore domains in the structure. In this
study, some weak intense shoulder peaks are observed on the broad Raman peak of Ln= Gd
11

and Dy, which indicates the presence of weakly ordered pyrochlore domains in the material.
Therefore it can be concluded from the vibrational studies of Ln2Zr1.5Hf0.5O7 that, the Ln = Pr
and Nd ceramics are perfect pyrochlores, Ln = Er is a perfect defect fluorite and Ln = Gd and
Dy are defect fluorites with weak pyrochlore domains. Previous studies on the Ln2Zr2O7
shows that Ln = Gd is the phase boundary between pyrochlore and fluorite phase boundary
[21]. The shift of phase boundary observed in this study is shifted to Ln = Dy, which is due to
the substitution of isovalent Hf4+ ions in the Zr4+ sites.
Fig. 3 TEM images of (a) Pr2Zr1.5Hf0.5O7, (b) Er2Zr1.5Hf0.5O7 and (c) their SAED patterns
12

The size and morphology of the prepared powders are analysed by using
Transmission Electron Microscopy (TEM). TEM images of Ln = Pr and Er are shown in Fig.
3 (a) and (b) respectively. It is clear from the images that, the nano ceramic particles are
highly agglomerated with less than 10 nm size. This smaller size of the particles are
comparable with the crystallite size calculated from the XRD. The high agglomeration
observed in the image is the characteristic nature of combustionally derived nano particles. A
further insight into the structure of the prepared nano particles are explored by using High
Resolution Transmission Electron Microscopy (HRTEM), which are shown in the inset of
Fig. 3(a) and (b). The HRTEM images shows the regularly ordered atomic arrangements in
the materials. The interplanar spacing measured from the HRTEM images shows that Ln = Pr
forms the pyrochlore structure whereas Ln = Er forms as a fluorite type structure. In order to
differentiate the pyrochlore and fluorite structure, a combined Selected Area Electron
Diffraction (SAED) pattern is created as shown in Fig. 3(c). The ring nature of the materials
indicates the polycrystalline nature of the fine nano powder. The slightly larger radius of the
Ln = Er as compared with the Ln = Pr indicates the smaller d spacing of the former with
respect to the latter, which is consistent with the X-ray diffraction analysis.
The UV-Visible diffuse reflectance spectroscopy (DRS) is a convenient method to
analyse the absorption behaviour and optical properties of ceramic materials. The UV-Visible
absorption spectra of the prepared nano ceramics between 200 and 700 nm are shown in Fig.
4(a). All the prepared lanthanide materials except Ln = Pr, shows high absorption in the UV
region and then exponentially decreased towards the visible region. This exponential decrease
is due to the defect states present in the forbidden region. However Ln = Pr shows a wide
range of absorption, which starts from the visible region and extends up to the UV region.
Balakrishna et al. reported that Pr3+ ions have several absorption bands near 400 nm
corresponding to the transitions from the ground level (3H4) to the higher energy levels of the
13

Pr3+ ions [22]. However the defects states present in the forbidden energy gap may widen the
absorption towards the forbidden gap region. A similar type of spectrum was reported in
previous literature [23,24]. The absorption band edge (wavelength at which absorption starts)
of the prepared nano ceramics can be identified by extrapolating the linear part of the
absorption spectrum to the wavelength axis. The measured absorption band edge values are
presented in Table 3. It is clear from the table that Ln2Zr1.5Hf0.5O7 (Ln= Nd, Sm, Gd, dy and
Er) nanoceramics shows the absorption band edge in the UV region, whereas Ln = Pr shows
an absorption edge at the visible region (647 nm).
Fig. 4(a). UV-visible Absorption spectra of Ln2Zr1.5Hf0.5O7 ceramics
14

Fig. 4(b). Tauc’s plots of Ln2Zr1.5Hf0.5O7 nanoceramics
Table 3 Absorption and emission properties of Ln2Zr1.5Hf0.5O7 nanoceramics
UV visible spectroscopy
Photoluminescence spectroscopy
Ln
Ceramics
Absorption
Band gap
Band edge
(eV)
Transition
Excitation
(nm)
Emissions
(nm)
levels
(nm)
Pr
647
3.71
5.01
414
324
620
1D2 → 3H4
Due to the
structure
defects
368
411
Nd
271
Sm
Gd
261
243
5.21
5.33
404
278
605
301
4G5/2 → 6H7/2
6P5/2 → 8S7/2
4F9/2 → 6H15/2
4F9/2 → 6H13/2
4S3/2 → 4I15/2
4F9/2 → 4I15/2
471
578
529
667
Dy
Er
258
246
5.43
5.47
386
445
The electronic transitions in the forbidden gap obeys the Tauc’s relation [25],
F(R).hʋ = A (hʋ – Eg) m
15

where A is a constant, Eg is the band gap energy, h is the Plancks‘s constant, F(R) is
the Kubelka-Munk function, ʋ is the frequency and m is an index, which assumes values
between 1/2 and 3 depending on the nature of the transition mechanism responsible for
absorption of photon in the energy gap. The exponent m = 1/2, 3/2, 2 and 3 are corresponding
to the direct allowed, direct forbidden, indirect allowed and indirect forbidden transitions
respectively. The value of Eg can be evaluated by extrapolating the linear part of [f(R).hʋ]1/m
versus hʋ graph to the point [f(R).hʋ]1/m = 0. Earlier studies on Ln2Zr2O7 ceramics showed
that these materials are direct band gap materials [21,26,27]. Therefore m= ½ is selected for
the present study. The Tauc’s plots for the prepared materials are shown in the Fig. 4(b) and
the corresponding estimated band gap energy are presented in Table 3. The measured band
gap values of these materials are not comparable with the energy values corresponding to the
absorption band edge wavelength. This may be due to the presence of defect states in the
band gap region. The band gap values reported in this study are much higher than the
previous reports. Uno et al. studied the band gap of Ln2Zr2O7 (Ln = La, Ce, Nd and Sm)
through the photochemical measurements and it was found that the band gap values of these
ceramics lie between 2.53 and 3.52 eV [27]. The reported band gap of Dy2Zr2O7 in the earlier
study was also less than the Ln = Dy in the present study [26]. Moreover the measured band
gap energy values in this study is higher than that of the previously reported band gap values
of Ln2Zr2O7 nanoceramics prepared by the same method [21]. That is the 25% of Hf ion
substitutions at Zr sites of Ln2Zr2O7 (Ln = Pr, Nd, Sm, Gd, dy and Er) considerably increases
the band gap energy of the materials.
The optical emission properties of the prepared nanomaterials are analysed by using
Photoluminescence spectroscopy (PL). Fig. 5 shows the PL emission spectra of the prepared
nanocreamics, the excitation wavelengths used for this emissions are included in Table 3. The
absorption spectra of Ln2Zr1.5Hf0.5O7 with Ln = Pr, Sm and Er shows several peaks in the
16

blue region. Therefore the PL emission spectra of these materials are monitored under the
blue region excitations, all the three materials shows orange red or red emissions. By exciting
under 414 nm, Ln = Pr shows the intense orange red emission (620 nm) of Pr3+, the 404 nm
excitation on Ln = Sm also shows an orange red emission (605 nm) of Sm3+ and the 445 nm
excitation on Ln = Er shows an intense red emission (667 nm) of Er3+. That is these materials
are potentially useful in blue-based white LEDs. The electronic transitions corresponding to
the emission spectra of these materials are also included in the Table 3, which agrees well
with the previous reports [22,28,29]. The emission spectra of Ln = Nd shows two low intense
broad emissions near 368 and 411 nm. Nd3+ ion is well known for its near infrared (NIR)
luminescence. Therefore the observed peaks cannot be simply attributed to the Nd3+
transitions levels. A. Zhang et al. reported that the PL emission of Nd3+ near 360 and 400 nm
is due to the structure defects in the nano crystals [30]. The PL emission of Ln = Gd shows
broad UV emission near 300 nm for an excitation of 278 nm. The observed spectrum is
consistent with previously reported data [31]. However the PL emission spectra of Ln = Dy
shows two peaks, one at the blue region (471 nm) and the other at the yellow region (578
4
6
nm). The F9/2 – H15/2 transition of Dy3+ ions corresponding to the 471 nm is magnetically
4
6
allowed, while the F9/2 – H13/2 transition corresponding to the 578 nm is a forced Electric
Dipole Transition (EDT). The intensity of EDT transition depends on the local environment
of the Dy3+ ions. When Dy3+ ions are located at low-symmetry sites with no inversion
4
6
centers, the F9/2 – H15/2 transition is predominant [32]. The dominancy of the 471 nm peak
in Ln = Dy in this study shows that Dy3+ ions are located at the low symmetric matrix
17

Fig. 5 PL emission spectra of Ln2Zr1.5Hf0.5O7 nanoceramics
The dielectric and ionic transport properties of the prepared materials can be studied
in dense ceramic materials. For this, the prepared nano materials are thoroughly mixed with
5wt.% of poly vinyl alcohol (PVA) in an agate mortar until a uniform mixing is attained.
These powder samples are then dried and again grinded in the agate mortar for half an hour.
18

Disc shaped pellets of 10 mm diameter and 2 mm thickness are then prepared from this
power materials under a pressure of 20 MPa. The green compacts are then sintered in a
o
resistive heated furnace at a constant heating rate of 5 C/min and soaking time of 90 min.
The sintering temperature and the density achieved by each ceramic material is tabulated in
Table 4. The prepared bulk ceramics achieves a relatively higher density at the lower
temperature as compared with the previous reports [21,33]. The homogeneous size
distribution of the nano particles may be the cause of higher density and reduced sintering
temperature in this study. Nano particles have a high surface to volume ratio and hence large
surface energy. The tendency to reduce this excessively large surface area per unit volume of
the nanocrystalline powders is the mechanism that drives the sintering process. Also the
sintering stress in the 10 nm particle is 100 times larger than that of the 1 μm particles [34].
That is the sintering temperature of Ln2Zr1.5Hf0.5O7 (Ln = Pr, Sm, Nd and Er) is considerably
decreased as compared with the sintering temperature of undoped Ln2Zr2O7 prepared by the
same preparation method [21]. Thus the partial substitution of Hf ions in Ln2Zr2O7 improves
the density and decreases the sintering temperature. However the sintering temperature of Ln
= Gd and Dy are slightly increased, which may be due to the improved density of these
materials as compared with the previous report. The density of Hf4+ ions is much larger than
that of Zr4+, since Hf ions have larger atomic mass and comparatively similar ionic radius
with the Zr ions. Therefore the presence of Hf ions may imposes additional stress to the
nearby particles, which in turn improves the sintering density.
The surface morphology of the densified ceramics, Ln = Nd, Sm, Gd and Dy are
examined by using field emission scanning electron microscopy (FESEM). The FESEM
image of the thermally etched pellet samples are shown in Fig. 6 (a). A highly compacted
grain structure with less porosity and well separated grain boundaries are visible in the
19

images. Some small black spots are observed on the FESEM images of Ln = Sm and Dy
ceramics, which is due to the diffused silver during the impedance analysis (discussed later).
Table 4 Sintering density of Ln2Zr1.5Hf0.5O7 ceramics
Sintering
temperature
(OC)
Experimental
density
Ln
ceramics
% of density
(g/cm3)
Pr
Nd
Sm
Gd
Dy
Er
1420
1400
1420
1485
1475
1470
6.29
6.75
6.99
7.18
7.36
7.85
93.54
99.53
99.22
98.32
96.05
99.29
It is clear from the images that the grain size of Ln = Gd and Dy are higher than that of the Ln
= Nd and Sm ceramics. The higher grain size of Ln = Gd and Dy are due to its higher
sintering temperature as compared with the others. The grain size distribution in these images
are collected by using imageJ software. A histogram is plotted by taking grain size along the
x-axis and grain counts along the y- axis as shown in Fig. 6 (b). The average grain size in
from this histogram data is measured by fitting Gaussian curve to this histogram. The average
grain size measured for the Ln = Nd, Sm, Gd and Dy ceramics are, respectively, 0.21 ± 0.06,
0.26± 0.07, 1.35 ± 0.26 and 0.84 ± 0.27 m. The standard deviation measurements of Ln =
Gd and Dy are higher than other material, which indicate that the particle size distribution of
these ceramics are scattered to large extent. According to classical sintering theories, during
the sintering process the grain boundary between two grains migrates form a larger grain to a
20

smaller grains due to the atomic diffusion [35]. Therefore the grain boundary diffusion and
the grain size of the ceramics gradually increases with increase of temperature.
Fig. 6 (a) SEM image of Ln2Zr1.5Hf0.5O7 ceramics (Ln = Nd, Sm, Gd and Dy)
Fig. 6 (b) Histogram of grain size distribution in Ln2Zr1.5Hf0.5O7 ceramics (Ln = Nd, Sm, Gd
and Dy)
21

Fig. 7(a). Cole-Cole plot of Sm2Zr1.5Hf0.5O7
Impedance spectroscopy can be used to determine the electrical properties of the
prepared dense ceramic materials. Fig. 7(a) shows the Cole-Cole plots of Ln = Sm ceramics
between 500 and 850 OC. The impedance plot at 500 OC is characterised by a single semi-
circular arc whose centre lies below the x axis, above 600 OC an inclined spike is appeared
along with the semi circles. The spike at the low frequency region on the right side of the plot
is due to the space charges between the pellet and the electrode and the arc at the high
frequency region on the left side of the plot is due to the combined effect of grain and grain
boundary [36]. Due to the instrument’s limitation, the semicircle corresponding to the grain
and grain boundary effect is not completely visible at high temperature. It can also be seen
that with the increase of temperature the size of the semicircle gradually reduces and shift
towards the origin. This indicates that the conduction process in the material is thermally
activated. Other prepared materials in this study (Ln = Nd, Gd, Dy, and Er) also shows a
similar impedance behaviour as that of Ln = Sm.
22

Fig. 7(b). Cole-Cole plot of Ln2Zr1.5Hf0.5O7 ceramics at different temperature.
Fig. 7(b) shows the variation of impedance with the substitution of Ln ions. It can be
seen that impedance of the Ln2Zr1.5Hf0.5O7 (Ln = Nd, Sm, Gd, Dy, and Er) gradually
decreases as the Ln ions changes from Nd to Dy and a sudden increase is observed for Ln =
Er. It is identified from the vibrational studies that the oxide ion disorder in the pyrochlore
structure increases as Ln3+ ion changes from Nd to Er with pyrochlore to defect fluorite
transition at Ln = Dy. That is the impedance of the Ln2Zr1.5Hf0.5O7 pyrochlore decreases with
23

the increase of disorder in the structure up to the pyrochlore-fluorite boundary and then
increases with the increase of disorder. This is due to the fact that the number of mobile
charge carriers increases with the increase of disorder and hence the impedance decreases.
However the large increase of impedance in Ln = Er, irrespective of the increase of disorder
is due to the increased ion- ion interaction in the material. Another interesting observation
from the Fig. 7(b) is that the impedance of the Ln = Er ceramic is three times larger than that
O
of the impedance of Ln = Nd ceramic at 500 C. This large difference in impedance is
gradually decreased with increase of temperature and it became comparable (but impedance
of Ln = Er is greater than Ln= Nd) at 750 OC. This behaviour of the Ln = Er ceramic indicate
that the number of mobile oxide ions liberated from the ion- ion interaction drastically
increases due to thermally activated conduction.
The total electrical conductivity of the bulk ceramic materials are calculated by using
the formula σ = G.t/A, where G is the measured conductance, t is the thickness and A is the
area of the ceramic material. The temperature dependence of electrical conductivity of the
dense materials can be studied by using Arrhenius relation
σ.T = σ0 exp(-Ea /KB T)
where the pre-exponential number σ0 is a measure of number of mobile charge
carriers in the material, Ea is the activation energy, KB is the Boltzmann constant and T is the
absolute temperature. Fig. 8 shows the Arrhenius plot of the ceramic materials, the straight
line fit to points on the plot shows that the ionic conduction of the materials is thermally
activated. The higher slope observed in the straight line fit of Ln = Er confirms the sharp
variation of impedance observed in the Cole – Cole plot. The values of σ0 and Ea can be
respectively evaluated from y-intercept and slope of the straight lines. Table 5 shows the
24

variation of activation energy and pre-exponential number with the variation in ionic radius
ratio r(Ln3+)/r(Bav4+). The average ionic radius of B site cation (Zr4+/ Hf4+) is calculated as
Bav4+ =0.75 r (Zr4+) + 0.25 r(Hf4+)
Díaz-Guillén reported that the activation energy of defect fluorites are less than that of the
pyroclore structures and their reported values are comparable with the present study [37].
Activation energy values of Ln2Zr2O7 reported by Yamamura et al. [9] are slightly less than
that of the activation energy values presented in the Table 5, which are be due to the smaller
grain size of the materials in the present study. However the variation of activation energy
with average cationic radius ratio in the present study is similar to the previous reports
[15,21].
Fig. 8. Arrhenius plot for Ln2Zr1.5Hf0.5O7 ceramics
25

Table 5 variation of ionic transport properties with Ln3+ ions
Conductivity
Pre-
exponential
number
×104
Radius ratio Activation
(Scm-1)
Ln
Ceramics
r(Ln3+)/r(Ba
energy
(eV)
4+
at 750 OC at 850 OC
)
v
×10-4
×10-3
1.55
1.50
1.47
1.43
1.40
1.09
1.07
0.89
0.88
1.18
3.62
6.41
1.11
1.31
10.30
1.49
0.49
Nd
Sm
Gd
Dy
Er
3.31
3.92
5.03
1.47
1.08
1.30
1.63
0.61
The A2B2O7 pyrochlores contains an intrinsic oxide ion vacancy at the 8b site. The
activation energy required for the oxide ion conduction in this pyrochlore structure is higher
than that of the proton conduction [38]. The activation energy values reported in this study
varies between 1.18 and 0.88 eV, which agrees with oxide ion migration in the pyrochlore
crystal lattice [38,39] Also, it is clear from the table that the decrease of radius ratio gradually
decreases the activation energy of Ln2Zr1.5Hf0.5O7 up to Ln = Dy and then increases for the
Ln = Er. That is the oxide ion movement in disordered fluorite structure requires more
activation energy than that of ordered pyrochlore structure. The σ0 value of the ceramics
shows an increase with the decrease of radius ratio except for Ln = Gd and Dy. That is the
number of mobile charge carriers present in the materials increases with decrease of radius
ratio, since the σ0 value is a measure of mobile charge carriers present in the sample. The
decrease of Ea will facilitate an easy path for the oxide ion hoping conduction and increase of
σ0 will increase number of mobile oxide ions in the crystal lattice. Both of the factors will try
improves the oxide ion conductivity of the materials. Thus conductivity of the
26

Ln2Zr1.5Hf0.5O7 ceramics gradually increases with the replacement Ln3+ ions from Ln = Nd to
Dy. Instead of increasing the ionic conductivity with the increase of oxide ion concentration,
Ln = Er ceramic material shows a decrease in its conductivity. This contrary behaviour of Ln
= Er ceramic is due to the ion – ion interaction and increase of activation energy. When the
number of mobile oxide ions increases, a weak ion – ion interaction is developed between the
oxide ions and it will restricts the hopping motion of charge carriers [40]. Thus the available
number of charge carriers participating in the conduction process will decreases. Therefore
the ionic conductivity of defect fluorite structured Ln = Er shows the lowest conductivity
value in the series. The maximum ionic conductivity observed at 800 OC in this study is 1.63
-1
×10-3 Scm for the Ln = Dy ceramic with radius ratio 1.43. Yamamura et al. reported that
the maximum oxide ion conductivity in the Ln2Zr2O7 is at the pyrochlore fluorite phase
boundary (Ln = Gd) with radius ratio 1.26, ionic radius values of the B site cations are in the
8 fold coordination [12]. By considering the 6 fold coordination for the B site cations, the
corresponding radius ratio value in this study is 1.46. Earlier conductivity studies on the
Ln2Hf2O7 series also show that Ln = Gd has the maximum conductivity with cationic radius
ratio 1.48 [41]. In this study the maximum conductivity is observed at the vicinity of
pyrochlore to fluorite transition (Ln = Dy) but the average cationic radius ratio of the phase
boundary is shifted to 1.43. That is the oxide ion conductivity of the prepared materials in the
present study is not only depending on the oxide ion disorder but also the preparation
methods. The measured conductivity values in this study are slightly less than that of the
previously reported conductivity values of Ln2Zr2O7 prepared by the same method [21]. This
slight decrease observed in the oxide ion conductivity is due to the reduced grain size of the
bulk ceramics as compared with the above reported study. The grain conductivity of bulk
ceramics usually decreases with the decrease of grain size of the materials. The structural
analysis performed in this study shows that Ln = Dy ceramic is neither an ordered pyrochlore
27

nor a disordered defect fluorite structures. It can be considered as a defect fluorite structure
with weakly ordered pyrochlore domains. That is the partial substitution of Hf4+ ions at the
Zr4+ sites shifts the order disorder transformation boundary and peak ionic conduction point
from Ln = Gd to Ln = Dy in Ln2Zr2O7.
Fig. 9 shows the fitted curve and the equivalent circuit of the Ln = Gd ceramic at 500 OC
Table 6 Equivalent circuit parameters for grain and grain boundary effects.
Resistance values at selected temperatures
Range of capacitance
500 OC
600 C
700 OC
O
Cermics
Cg ×10-11
F
Cgb ×10-10
F
Rg
Ω
Rgb
Ω
Rg
Ω
Rgb
Ω
Rg
Ω
Rgb
Ω
Nd
Sm
Gd
Dy
Er
4.91 – 7.35
0.82 – 8.17
10.8 – 14.0
12.1 – 18.9
3.35 – 6.82
4.84 – 5.56
3.41 – 7.20
2.10– 5.19
4.58 – 9.17
1.0 – 2.7
81628
28273
15403
13030
54513
25350
6923
10821
4368
3129
2196
19343
7999
3745
1525
1283
17900
2548
926
1209
847
618
390
893
679
5426
577
295610
183910
4041
28

The grain, grain boundary and electrode contribution to the electrical properties of the
bulk ceramic material are further analysed by fitting suitable circuit elements. In this study
the Cole-Cole plot of the Ln2Zr1.5Hf0.5O7 shows a depressed semicircle with an inclined spike.
The depressed semicircle at the intermediate and high frequency region is due to the
combined effect of grain and grain boundary, it can be fitted with a series combination of two
parallel resistance – capacitance elements. Fig. 9 shows the fitted curve and the equivalent
O
circuit of the Ln = Gd ceramic at 500 C. In this study, Z-view 3.2 software is used for the
curve fitting analysis. The calculated values of grain resistance (Rg), grain boundary
resistance (Rgb), grain capacitance (Cg) and grain boundary capacitance (Cgb) for selected
temperatures are shown in the Table 6. It is clear from the Table 6 that Rg is higher than that
of Rgb at all temperatures and chemical compositions. This may be due to the smaller grain
size of the bulk ceramic material. The grain resistance of a material usually increases with
decrease of grain size. It is also clear from the table that the Rg and Rgb values decreases with
increase of temperature in accordance with the Arrhenius law of thermally activated ionic
conduction. Fig. 10 shows the Arrhenius plot for the grain conduction. The linear behaviour
of the Arrhenius plot shows that the grain conduction is thermally activated. Variation of the
activation energy (Eag) and pre-exponential values (σ0g) of the Ln2Zr1.5Hf0.5O7 grain
conduction is graphically shown in the Fig. 11 The Eag and σ0g values of the grain conduction
of Ln2Zr1.5Hf0.5O7 shows a similar behaviour as that of the total electrical conductivity
(discussed earlier). That is, the Eag value decreases gradually with the decrease of
r(Ln3+)/r(Bav4+) ratio, reaches a minimum for Ln = Dy and then increases with the decrease of
radius ratio. The capacitance values evaluated for the material are shown in Table 6. From
this table it can be identified that Cgb is always greater than the Cg value of the bulk ceramics.
The Cg and Cgb values are of the order of 10–11 and 10-10 F, respectively, which are
comparable with the grain and grain boundary capacitance in the previous studies [14,42].
29