CHAPTER I

INTRODUCTION

1.1 A Brief History of Fast Ionic Conductors

Solid state ionics concerns properties associated with the motion of ions in

solids [1]. It provides scientific support for a wide variety of advanced

electrochemical devices such as batteries, fuel cells, chemical sensors, gas separation

membranes and ionic switches. Numerous solids with high ionic conductivity at room

temperature are extensively used in technological applications and are named as fast

ionic conductors (FIC). One of the indispensable uses of FICs is as electrolytes in

battery applications; hence those are referred as solid electrolytes. Longer life, high

energy density, less possibilities of leak etc., are some of the advantages of using solid

electrolytes in electrochemical devices against the liquid electrolytes.

The study of ionic conduction in solids was initiated in 1838, when Faraday

discovered PbF2 and Ag2S as good conductors of electricity. In subsequent years, a

variety of solids exhibiting appreciably high ion conductivity at their operating

temperature was recognized. Silver ion conducting solids of the type, RbAg4I5, were

discovered in early sixty and were used in electrochemical cells [2]. The Na-β-

alumina was successfully used in Na/S batteries by Kummer and Weber [3]. Hong [4]

and Goodenough et al., [5] reported in 1976 a high ion conducting skeleton structure

having polyhedral units popularly known as NASICON (Na Super Ion CONductor).

The polyhedral skeleton structure consists of rigid (immobile) subarray which

provides a large number of three dimensionally connected interstitial sites suitable for

long range motion of monovalent cations. In contrast to β-Al2O3,3 in this mixed

zirconate-phosphate compound, Na+-ions move within three dimensional channels of

the structure. Analogous compounds, such as LISICON, were synthesized shortly

thereafter [5]. From 1970 onwards, a number of studies focusing on the synthesis and

characterisation of lithium ion conductors were reported. The enhanced interest on

lithium ion conductors were driven by the small ionic radii, lower weight, ease of

handling and its potential use in high energy density batteries. Li4SiO4 as well as its

non-stoichiometric solid solution Li4-3xAl3SiO4 (0≤x≤0.5) were promising Li fast ionic

conductors.

2

Otto et al., [6], was the first to report the glass compositions (M2O-SiO2-Bi2O3

with M=Li, Na) with exceptionally high alkali-ion conductivity at relatively low

temperatures, thereby demonstrating that long range order was unnecessary to support

the fast ion conduction. Further, important developments in solid electrolytes in this

period included the discovery of fast ionic conduction in inorganic glasses [6] and

polymers [7].

The general requirements for practical solid electrolytes are: (1) high ionic

conductivity (2) negligible electronic conductivity (3) stability with respect to

adjacent phases, thermal and electrochemical decomposition (4) suitable mechanical

properties (5) ready availability of chemical constituents (6) ease of fabrication (7)

reasonable cost and are particularly suitable in compact power batteries used in pace-

makers, mobile telephones and laptops.

Similar to electronics, progress in solid state ionics motivated major

technological developments, especially in the field of energy storage and conversion.

The mixed conductors-electronic and ionic-play a critical role, particularly as

electrodes in electrochemical devices. Future developments in the industrialized world

depend largely on the existence of sufficient energy resources. The consumption of

fossil fuels, however, depletes available reserves and leads to severe environmental

problems, including acid rain and global warming by emission of NOx, SO2 and CO2.

Methods for decreasing these emissions include improving the efficiency of

combustion engines or by converting the hydrocarbons, CO and NOx emissions into

environmentally harmless products. Gas sensors operating in the feedback mode with

microprocessors and fuel injectors are essential for such tasks. Indeed, one of the

major successes of solid state ionics in this century is the use of sensors based on

ZrO2 solid electrolyte in almost every automobile produced worldwide. One of the

major challenges in solid state ionics continues to be the development of more

selective electrochemical sensors capable of detecting a broader variety of gases,

including NOx, SO2, CO and CO2 with response times often in the millisecond range.

Besides energy conversion and storage, solid state ionics offers potential

solutions for energy conservation. This is based on the development of smart

electrochromic windows in which charge transferred between anode and cathode of

an appropriately designed thin film battery, applied to a glass window can be used to

vary the optical reflectivity of the multilayer film over wide limits. Thus, during

midday, when solar absorption is normally highest, a high reflectivity can be tuned in

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order to minimize the heat absorption, thereby decreasing the need for air

conditioning. Alternatively, on cold days, the reflectivity edge can be shifted to longer

wavelengths, thereby transmitting the visible spectrum, to trap the heat. Remaining

challenges include means for preparation of uniform coatings over large area at low

cost and continued efficient operation for long periods.

Nanocrystalline materials have shown several interesting properties and are

explored in different applications [8]. However, in many cases, the physical

mechanisms responsible for the interesting properties are not completely known yet.

One important example is the enhancement of ionic conductivity in nanostructured

solid electrolytes. In fact, now a days nanoionics is one of the recent fields of research

related to nanomaterials and several applications are being explored, e.g. rechargeable

lithium ion batteries, gas sensors and solid oxide fuel cells. The materials in nanoscale

have a large number of grain-boundaries and atoms present on the surface. These

surface atoms influence the defect thermodynamics considerably.

1.2 Classifications of Ionic Conductors

Electrical conduction occurs by long range migration of either electrons or

ions. Usually conduction by one or other type of charge carrier dominates, but in

some inorganic materials, both ionic and electronic conductions are appreciable. The

specific conductivity for any charge carrier is:

𝜎 = ∑ 𝑛𝑖𝑒𝑖𝜇𝑖𝑖 (1.1)

where, 𝑛𝑖 is the number of charge carriers of the species i, 𝑒𝑖 is its charge and 𝜇𝑖 is its

mobility. For few of the materials other than metals, conductivity increases with

temperature.

Migration of ions does not occur to appreciable extent in most of the ionic and

covalent solids. Atoms tend to be essentially fixed on their lattice sites and can move

via crystal defects. Only at high temperatures the defect concentrations become quite

large and conductivity becomes appreciable. The conductivity of NaCl at 800ºC, i.e.,

just below its melting point is 10-3Scm-1, whereas, at room temperature pure NaCl is

an insulator [9].

There exists a small group of solids namely solid electrolytes or fast ionic

conductors, in which, one set of ions can move easily. Such materials have special

crystal structures, i.e., there are open tunnels or layers through which the mobile ions

4

can move easily. The conductivity values of Na+ ion migration in β-alumina at 25ºC is

10-3Scm-1, is comparable to those observed for strong liquid electrolytes. Existence of

disorder or defect is necessary for ionic transport in solids; and the density of defects

depends on various factors like structure, temperature, presence of impurity ions,

nature of chemical bonding between the constituent ions etc. [10]. A lot of research

work is being carried out in order to find the materials with improved ionic

conductivity. The ion conducting materials are classified based on doping approach

and structure.

1.2.1 Doping Strategies

Homogeneous doping with appropriate elements is one approach to obtain

materials with high ionic conductivity. These types of compounds contain an ion with

different valency to maintain the bulk electrical neutrality and the supplementary

defects are created. Depending on the types of defects different materials are

investigated.

1.2.1.1 Fluorite-Type Compounds

Oxides with fluorite-type structure (e.g. ZrO2, CeO2, ThO2) can form solid

solutions with lower valence cations like CaO, Y2O3 and other rare earth oxides

resulting in 10%-15% oxygen deficiencies. These materials find application in

cathode materials, gas sensors, catalysts and in oxygen storage materials. Fluoride ion

conduction is common in materials with fluorite structure; where PbF2 shows

significant conductivity, but CaF2, SrF2 and BaF2 are moderate fluoride ion

conductors [11]. Aliovalent doping is used to enhance the low temperature ionic

conductivity in these materials [12].

1.2.1.2 Perovskites

Compounds in perovskite or related structures have many applications in high

performance electroceramics, including high-TC superconductors, dielectrics,

piezoelectrics and ferroelectrics [13]. The chemical formula of these compounds is

ABO3, few of these materials show mixed conductivity. It can accommodate dopant

cations in A and B sites, which improve its electronic or oxygen ion conductivity [14].

Depending on the cation radii, ideal cubic structure may deform to lower symmetry.

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(i) Solid Oxide Fuel Cell Electrode Materials

Oxygen ion conduction in perovskite like La1–xCaxAlO3 and CaTi1–xAlxO3 is

3x10-2Scm-1 at 1000°C. The ability of perovskite structures to accommodate

multivalent titanium ion contributes to the mixed ionic and electronic conductivity.

Goodenough et al.,[15] examined highly oxygen deficient perovskite related

brownmillerite structures of the form A2B2O5, at elevated temperatures that change to

perovskite structure.

(ii) Proton Conductors

Most of the proton conductors discovered initially were hydrates and were

thermally unstable [16]. Iwahara and coworkers [17] did pioneering investigations on

proton conduction at high temperatures in perovskite oxides, ABO3, where A=Ba, Sr

and B=Ce, Zr.

1.2.2 Structurally Disordered Materials

1.2.2.1 Crystalline Compounds with Disordered Sub-Lattice

In this type of materials, highly disordered sub-lattice exists, where the

number of equivalent sites is greater than the number of ions to fill it [18]. This results

in exceptionally high mobile carrier concentrations (1022cm–3) and diffusion

coefficients are comparable to those in liquids (typically ~10-5cm2s-1). The carrier

concentrations are much higher than in classical solid ionic conductors (typically 1015

to 1018cm-3) where, the causes of point defects are thermal activation, deviations from

stoichiometry or doping.

(i) Noble metal Halides and Chalcogenides

High ionic conductivity and low activation energy of α-AgI is due to the

presence of many randomly accessible almost equivalent sites with small energy

barriers in the body centered cubic structure. The α-CuBr is structurally analogous to

α-AgI, exhibits high ionic conductivity. Major work was done to stabilize the α-

structure to the room temperature.

(ii) Bismuth Oxide and BIMEVOX Compounds

The Bi2O3 and mixed bismuth oxides exhibit high oxide ion conductivity. The

low temperature semiconducting α-modification transforms to δ-phase at 730˚C [19]

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and is accompanied by an oxide ion conductivity jump of almost three orders of

magnitude. δ-Bi2O3 presents disordered fluorite type structure, where a quarter of the

oxygen sites are intrinsically empty. The polarizability of bismuth cation is

responsible for its high anion conductivity and is one to two orders of magnitude

higher than doped ZrO2. Attempts to stabilize δ-phase to the lower temperature, led to

a new family called BIMEVOX compounds [20], in which vanadium is partially

substituted by low valence cations like Cu, Ni or Co. These compounds present high

mixed conductivity with remarkable oxygen ion conductivity at moderate

temperatures. These are promising compounds for use in oxygen separation

membranes with high permeation rates [21].

(iii) Pyrochlores

These compounds have the general formula, A(III)2B(IV)2O6X, where X is an

anion that can be oxygen. It is a superstructure of defect fluorite lattice with exactly

twice the lattice parameters. This material’s degree of disorder can be varied almost

continuously from low to high values using solid solution systems. Two types of

disorders are important in pyrochlores i.e., Frenkel like disorder on anion sub-lattice

and anti-site disorder on cation sub-lattice where A and B cations switch positions

[22].

(iv) NASICON Type Compounds

NASICON is a family of compounds of general formula, Na1+xZr2(P1–xSixO4)3

(0<x<3), where excess charge of the dopant is counter balanced by sodium ions. The

NASICON frame-work is built of (Si,P)O4, tetrahedra and ZrO6, octahedra, which

provide relatively open three dimensional network of sites for Na+ ions. This material

presents an unusually high ionic conductivity (optimized at x=2) and relatively low

activation energy. G. Govindaraj et al., studied the ac electrical properties of glassy

NASICON materials in detail [23]. NASICON has been proposed as the sensitive

membrane of Na+ ion selective electrode with better selectivity than classical glass

electrodes [24].

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1.2.2.2 Layered Compounds

(i) Lithium Insertion Compounds

Lithium insertion is essential for the development of rechargeable lithium

batteries. The first layered compound with Li+ ion conductivity is Li3N. Doping with

hydrogen, Li3-xHxN, is essential because the substitutional hydrogen leaves vacancies

for Li+ ion conduction. However, low decomposition voltage limits its use as solid

electrolyte in batteries.

(ii) β-Alumina

Crystal structure of this compound is related to that of Al2O3. In β-Al2O3

structure, planes with sodium ions are intercalated between blocks with the normal

spinel structure, hence a general formula NaAl11O17 is obtained. However, this

compound is not stable and can be stabilized by partial substitution of Al by divalent

cations like Mg giving β"-alumina. Sodium ions are mobile in the planes producing

anisotropic ionic conductivity in β-alumina [25]. The conductivity depends on ionic

radius and is best for ion size intermediate between sodium and potassium.

1.2.2.3 Tunnel Compounds

Tunnel compounds support ionic conduction along one dimensional channel.

Basic limitations are due to the anisotropy of ionic conduction. Examples are

hollandites, with general composition BaxMn8O16(x<2), with tunnels parallel to the

crystallographic c-axis. The priderites have the chemical formula, Ax(B,Ti)8O16

(A=alkali metal, B=Mg, Zr and Ga), which have the same structure and show alkali

ion conduction along the tunnels. Higher conductivities at high alternating current

frequency may relay on alkali ion oscillation between bottlenecks in the tunnels.

1.2.2.4 Amorphous Materials

Amorphous materials intrinsically exhibit three dimensional disorders. The

transport mechanisms in glasses are substantially less understood than in crystalline

materials. Advantages of glassy fast ionic conductors over crystalline include

isotropic conduction, absence of highly resistive and corrosion sensitive grain

boundaries, wide compositional flexibility and ease of fabrication. Fast ionic

conductivity is observed in many glasses, especially those with small cations like Ag,

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Cu, Li and Na [26]. Ternary glasses contain a network former (e.g. SiO2, B2O3, P2O5

and GeS2), network modifier (e.g. Ag2O, Li2O, Cu2O and Ag2S) and dopant

compound, mostly a halide (e.g. AgI, CuI and LiCl). As the term implies, the glass

network structure and its physical and chemical properties can be substantially

modified by addition of modifier. Dopant salts apparently do not interact strongly

with network; cation and anion dissolve into interstices of the glass structure.

Increased ion conductivity can be resulted from the carrier density and/or carrier

mobility. According to the weak electrolyte model, enhanced conduction is primarily

due to dissociation between alkali and halide ions. Others believe that the additives

induce major changes in the glass structure, and it has a significant effect on ion

mobility. In this model, a large fraction of carriers are assumed to be non associated

and free to move in the glass.

1.2.2.5 Composite Solid Electrolytes

The development of composite solid electrolytes gives a new degree of

freedom in solid state ionics. Mixtures can provide enhanced conductivity together

with other useful characteristics such as improved mechanical strength. The

conductivity enhancement in numerous ceramic composites was reported. Besides

Al2O3, other oxides such as MgO, SiO2, CeO2, TiO2 and ferroelectric BaTiO3 was

found to be the effective second phases for ionic conductivity improvement. More

recently, the composite effect was also observed in anion conductors, such as lead or

calcium fluoride and even in inorganic solids with trivalent cation conductivity, such

as aluminium and rare-earth wolframates [27].

1.3 Technological Applications of Fast Ion Conducting Materials

1.3.1 Batteries

Power storage devices require high energy density batteries. The highest

possible energy density is achieved using reactants with high free energy of reaction

and low mass such as lithium or sodium. This requires that the solid electrolytes

remain stable under highly reducing or highly oxidizing conditions. Silver and copper

fast ion conductors were initially used to design advanced solid state batteries. The

anisotropic conductivity of layered compounds led to the development of improved

solid electrolytes such as NASICON. There is an increasing demand for

microbatteries compatible with microelectronics technology related to the

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development of laptop computers or portable telephones. This led to the development

of a high energy density and long cycle life rechargeable lithium batteries. Suitable

solid electrolytes for lithium batteries should have the requirements: (1) high Li+ ion

conductivity (2) compatibility with lithium metal or high lithium activity electrodes

(3) high decomposition voltage and (4) ease of thin film fabrication.

Many systems based on metallic lithium suffered from problems due to metal

oxidation and poor rechargeability due to the formation of metallic dendrites. The

alternative is based on lithium insertion compounds LixWO2 and LiyTiS2, but was

unable to provide sufficiently high energy density. Improved rechargeable lithium ion

batteries were based instead on nongraphitic carbons as the lithium insertion anodes

[28]. The insertion of lithium into normal graphite was unsuccessful and other

materials tested include certain zinc blende structure materials. The successful

association of hard carbon insertion anodes with the high voltage LiCoO2 insertion

cathode was a major breakthrough [28]. Because of the relatively high cost of cobalt,

alternative systems based on other cathode materials such as LiNiO2 or LiMn2O4 are

currently under investigation. The three dimensional spinel LiMn2O4 is

environmentally friendly and inexpensive, but it suffers from several drawbacks,

especially storage losses [29]. Presently, electrolytes are nonaqueous liquids, such as

ethylene carbonate or propylene carbonate with dissolved metal salts, including

LiPF6. Hybrid electrolytes based on polymers impregnated with liquid electrolytes is

in the development stage [29].

1.3.2 Solid Oxide Fuel Cells

Solid oxide fuel cells provide many advantages over traditional energy

conversion systems. It includes high energy conversion efficiency, fuel flexibility

because of internal reforming, very low levels of NOx, SOx emissions and long

lifetime because of the elimination of corrosive liquids [30]. Operation at an elevated

temperature is necessary because the solid electrolytes have relatively low ionic

conductivities and slow electrode processes at low temperatures. Furthermore,

mechanical stability and compatibility of the materials require improvement and costs

need to be decreased. It consists of two porous electrodes separated by a dense

oxygen ion electrolyte. Several important material requirements must be satisfied for

the optimal operation of solid oxide fuel cells: (1) suitable electrical properties (high

oxygen ion conductivity and low electronic conductivity for electrolyte material and

10

mixed conduction for electrodes) (2) adequate chemical and structural stability under

conditions of operation (3) minimal reactivity and inter diffusion between components

and (4) matching thermal expansion coefficients.

1.3.3 Sensors

Monitoring of chemical processes in real time enables closer quality control of

the products. Electrochemical microsensors transform a chemical signal into an

electrical signal that is easy to measure, monitor and process. Ionic and mixed

conducting solids are basic materials for its development because they can be easily

miniaturized, for instance, in thin film form and can be operated at elevated

temperatures or in an aggressive environment. Sensor development is largely driven

by the following requirements: high sensitivity, selectivity, short response time,

reproducibility, long term stability, lifetime, reversibility, size and low cost. The three

major types of electrochemical sensors are [31]:

(i) Potentiometric Sensors

In this type, the concentration or partial pressure of a species is determined by

measuring emf of solid electrolyte. The potentiometric sensors are categorized in to

three groups: (a) Oxygen Sensors: The most successful commercial sensor is the

oxygen sensor using stabilized ZrO2 as solid electrolyte. (b) Gas Sensors: These

sensors detect the atmospheric pollution gases using cation conducting metal salts as

solid electrolytes like Ba(NO3)2 for NO2, K2SO4 for SO2, and K2CO3 for CO2. (c) Ion

Sensors: An ion sensor (selective electrode) measures ion concentrations in solution.

It contains an ion conducting membrane in contact with solution, an internal reference

and a metal connector. The measured membrane potential depends on ion activity

according to Nernst’s law. Sodium and lithium ions have been detected with

NASICON type membranes.

(ii) Amperometric Sensors The amount of excess oxygen or fuel in an exhaust can be determined by

measuring the limiting electrochemical current of amperometric oxygen sensors.

(iii) Conductometric Sensors

It uses the dependence of electrical conductivity of a material on the chemical

potential of gaseous species.

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1.3.4 Electrochromic Windows

Electrochromic light transmission modulators called smart windows, using

solid ionic conductors, are significant in energy savings by regulation of thermal

insulation. In such a system, the window is maintained transparent in the visible and

reflecting in the infrared during the winter, allowing penetration of sunshine but

blocking loss of interior heat. On the other hand, the window is rendered partially

opaque during hot summer days, decreasing the amount of radiations entering the

building. The electrochromic elements that color or bleach on insertion/deinsertion of

lithium or hydrogen ions, are sandwiched between two transparent thin film

electrodes and separated by solid electrolyte [32]. The transparent electrodes are

generally indium tin dioxide. Glass electrolytes appear as promising choices.

1.4 Polyanionic Frame-work

Polyanion compounds containing compact tetrahedral ‘anion’ structural units

(XO4)n- (X=P, Mo, W and S) with strong covalent bonding[33]. These units

networked to produce higher coordination sites such as oxygen octahedral, that are

occupied by other metal ions. Among the family of (PnO3n+1)(n+2)- polyanions, the

polyphosphates can be classified into: (1) orthophosphates (n=1) characterized by

(PO4)3- isolated units (2) pyro or diphosphates (n=2) in which P2O7 groups are formed

by corner sharing of two (PO4)3- units and (3) triphosphates (n=3) where three (PO4)

3-

units from (P3O10)5- anions. The polyanionic compounds considered under the present

study are NASICON and olivine type orthophosphates, which show great interest in

rechargeable batteries.

1.5 Review of NASICON

1.5.1 Structure of NASICON Polymorphs

Phosphates having NASICON structure are unique due to their properties

determined by structure and composition. These properties include high ion

conductivity, low (sometimes, negative) thermal expansion coefficient and low

thermal conductivity. Its hardness, chemical and thermal stability are also significant.

In addition, NASICON phosphates have catalytic activity or luminescent properties

and most attractive property is its high ionic conductivity. The general form of

NASICON (Na Super Ionic CONductor) type materials is Na3Zr2(SiO4)2(PO4) [34].

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The mixed frame-work, [T2(XO4)3]n- has an important role in crystal chemistry

of compounds with tetrahedral anions. In this frame-work, all oxygen atoms are

shared between XO4 tetrahedra and TO6 octahedra. The element T can be in different

oxidation states like +2, +3, +4, +5 and X=Si, P, As, S, Mo and W. Usually

monovalent cations compromise negative charges in the frame-work; even

multivalent cations are also suitable [35]. In the general, these compounds exist in

different polymorphs i.e., rhombohedral, orthorhombic, monoclinic and triclinic with

different space groups R3�c, P21/n and Cc. Figs. 1.1(a)-(d) show structural polymorphs

of NASICON.

Figure 1.1: Structure of NASICON polymorphs: (a) Rhombohedral symmetr (R3�c) (b) Orthorhombic (Pbna) (c) Monoclinic (P21/c) and (d) Triclinic (C1� ) symmetry. (Courtesy: J. Gopalakrishnan et al., and N. Anantharamulu et al.,)

In rhombohedral symmetry with R3�c space group, two different sites are

available for monovalent alkali atoms i.e., M1 sites surrounded by six oxygen atoms

and located at an inversion centre and M2 sites surrounded by ten oxygens and

disposed symmetrically around the ternary axis. If the octahedral M1(6b) site is fully

occupied, tenfold coordinated M2(18e) site is empty. In Fig. 1.1(a) red circles denote

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M1 and M2 positions in the interconnected channels within the frame-work. A pair of

near lying octahedra is linked through PO4 tetrahedra. Both these sites are arranged in

alternating way along the conducting channels. LiGe2(PO4)3 and LiTi2(PO4)3 are

crystallized in this symmetry and privileged occupancy of M1 site is deduced from

the neutron diffraction experiments [36]. Particular attention has been given to

titanium based materials having NASICON structure like Li1+xTi2-xNx(PO4)3, LiTi2-

xMx(PO4)3 and Li1-xTi2-xRx(PO4)3, where, N, M and R stand for tri, tetra and

pentavalent cations respectively [37]. Various methods are used to increase the

conductivity of parent LiTi2(PO4)3 for practical applications, which include addition

of Li2O or Li4P2O7 as a binder to increase the density of pellets [38]. Even though,

high energy ball-milling is widely used as a method for preparation of samples, few

literatures are available about the synthesis of nanocrystalline NASICON materials by

this technique. Attempts have been made to increase the conductivity of parent

materials by substitution of P in the frame by V or Sb [39]. The ball-milling technique

is used to enhance the electrical properties in LiTi2(PO4)3 and in Al and V substituted

LiTi2(PO4)3 materials.

The crystal structure and conductivity are determined by size of guest (or

mobile) cation in the frame-work. The MIIZr2(PO4)3 (MII= Mg, Ca, Sr, Ba, Mn, Co,

Ni, Zn, Cd and Pb) crystallize in NASICON or β-Fe2(SO4)3 type structure depending

on the ionic radii of M2+ [40]. Mg, Co, Ni and Zn are crystallized in β-Fe2(SO4)3 type

structure and showed an order-disorder transition between 600˚C-720˚C. But, Ca, Sr,

Ba, Cd and Pb compounds crystallize in NASICON type structure and the phase is

stable from room temperature to 1000˚C. The Mn based compounds sinter at 900˚C

crystallize in β-Fe2(SO4)3 type structure and those sinter above 900˚C crystallize in

NASICON structure. LiZr2(PO4)3 shows change in crystal structure with sintering

temperature; those sinter between 1000˚C-1100˚C is crystallized as mixture of two

phases [41].

The Li3Fe2(PO4)3 synthesized by Masquelier et al., [42] is crystallized in three

different forms, i.e., α-monoclinic (P21/n), β-monoclinic (P21/n) and γ-rhombohedral

(R3�c). Li3Fe2(PO4)3 is crystallized in Fe2(SO4)3 form of monoclinic symmetry (P21/n)

by solid-state reaction. The monoclinic cell results from small distortion of the

orthorhombic frame-work due to lithium ion ordering at room temperature. d’Yvoire

et al., and Bykov et al., [43] showed that monoclinic Li3Fe2(PO4)3 transforms

14

reversibly to orthorhombic structure on heating above 270˚C due to progressive

breaking of long range ordering of Li+ ions in the interstitial space. Na3Fe2(PO4)3 has

monoclinic C2/c structure with two formula units, i.e., Z=2 at room temperature [44].

The monoclinic β-phase is formed (C2/c) at 368K and above 418K rhombohedral γ-

phase exist (R3�c).

Ion exchange method has been explored for synthesis of NASICON

compounds that cannot be prepared by conventional techniques. The rhombohedral

form of Li3Fe2(PO4)3 is prepared by ion exchange method from Na3Fe2(PO4)3. The

room temperature α-Na3Fe2(PO4)3 has an ordered distribution of Na+ ions, that

resulted in a slight distortion of rhombohedral frame-work to monoclinic symmetry

(C2/c) [45]. The substitution of Na+ by Li+ through ion exchange leads to disordered

rhombohedral (R3�c) β-Li3Fe2(PO4)3. This technique is used to prepare Li3Fe2(AsO4)3

from Na3Fe2(AsO4)3, and the material is crystallized in rhombohedral symmetry [42].

Similar observations are reported by Delmas et al., [46] for ATi2(PO4)3 (A=Li, Na),

attributed to vacant M1 sites along C axis, resulting in higher electrostatic repulsion

between TiO6 octahedra.

In general, an insignificant distortion of hexagonal lattice reduces the

symmetry to monoclinic with P21/n or C2/c symmetries with one angle close to 120˚.

The monoclinic distortion (C2/c or Cc) splits M2 sites to: 3(M2)hex→(M3+2M4)mon and

octahedral M1 site shares faces with two adjacent TO6 octahedra. The formed groups

O3TO3M1O3TO3 are mutually linked along the hexagonal c axis by XO4 tetrahedra.

The infinite ribbons resulting from the linking are connected together normally to c

axis by XO4 tetrahedra and M2 sites are situated between these ribbons. In

rhombohedral phase, building blocks are aligned parallel to hexagonal c-axis, while in

monoclinic phase; these are aligned perpendicular to one another.

In some cases, the symmetry is deteriorated to triclinic, as reported by M. Catti

et al., in LiZr2(PO4)3 [47]. In this case, unit cell may be chosen with one angle close to

120˚ and the other two close to 90˚. Thus, the structural resemblance to the above

mentioned phases is quite evident. The rhombohedral and monoclinic structural

modifications differ in three dimensional connectivity between T2(XO4)3 lantern units.

There is a significant difference between monovalent cation sites between two phases.

In orthorhombic symmetry (Pcan), cations fully occupy three distinct crystallographic

sites; M2 and M3 are in five fold coordination and M1 is in four fold coordination.

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1.5.2 Phase Transition in NASICON Polymorphs

The Na3Fe2(PO4)3 showed phase transitions at 368K and 418K [43]. The

hysteresis is observed in transition temperatures, i.e., there is a shift in the transition

temperatures while heating and cooling. The low temperature phase has weak

monoclinic distortion induced by interstitial occupancy of sodium ions and it’s

ordering [48]. The cation order disappears upon β→γ transition. In the case of

Na3Cr2(PO4)3 material, calorimetry experiments showed the phase transitions at 348K,

411K and 439K [49]. X-ray diffraction indicates that the unit cell symmetry is

improved upon these transitions and structural aspects of the transitions are discussed

by Genkina et al., [50].

The Na3Zr2PSi2O12 phase undergoes a single monoclinic to rhombohedral

phase transition unlike the above mentioned compounds [51]. However, this transition

occurs over a wide temperature range of about 70K. For lithium double phosphates,

phase transition data are less abundant. Monoclinic to rhombohedral phase transition

is observed in LiZr2(PO4)3 and transition temperature depends on annealing [52].

Later, another phase was discovered [53]; monoclinic to rhombohedral phase

transition in that phase was observed only at 573K. The phase transitions in sodium

scandium phosphate Na3Sc2(PO4)3 are studied in detail by L. Boehm et al., [54].

1.5.3 Conductivity in NASICON Polymorphs

The maximum conductivity of Na3Zr2Si2PO12 composition falls in the region

0.06 to 0.036Scm-1. The activation energy for conduction in this composition is 17-

30kJmol-1 and the best conductivity of NASICON family reached 10-1Scm-1 at 570K

[55]. The spread in activation energy is due to existence of minor phases, different

densities of samples and different techniques used for synthesis. The Na3Sc2(PO4)3

showed dc conductivity of 3x10-6 to 5x10-3Scm-1 at 300K. Activation energy for

conduction in β-phase is less compared to α-phase because more energy is required to

transfer cation between M1 and M2 sites [56]. The sodium ion has to pass through a

triangular face of trigonal anti-prism and size of face changes in course of phase

transition or with temperature. The key parameter that determines anti-prism size is

the nature of T and alkali ions. The indium analogue has even lower conductivity at

almost all temperature ranges but slightly higher than iron and chromium compounds

[57]. NaZr2(PO4)3 shows still lower conductivity even though its unit cell volume is

16

larger than scandium compounds [58]. In this material M1 site is fully occupied and

M2 sites are vacant; hence mobile charge carrier concentration is lower in this

material compared to Na3A2(PO4)3 type materials, where only two-third of M2

positions are occupied. While, LiZr2(PO4)3 is having higher conductivity in spite of its

lower cell volume compared to Na analogue. It is because of the smaller size of Li+

ions, it can overcome the narrow triangular faces of coordination polyhedra easily

[59]. Lithium phosphates of trivalent elements like Li3Sc2(PO4)3, Li3Cr2(PO4)3 and

Li3Fe2(PO4)3 are in orthorhombic structure in the high temperature phases [60].

Activation energy for conduction decreases for α to γ-phase transition; hence

conductivity is more at high temperature phases. But d’voire et al., did not find this

tendency [43].

Neutron diffraction is a valuable tool to determine the structural information

on lithium ion transport. Powder neutron diffraction of Li1+xFexTi2-x(PO4)3 with x=0.5,

at room and high temperatures give the complex pattern of lithium disorder, involving

sites lying within both M1and M2 cavities. This suggests high mobility of Li+ ions, in

contrast with results obtained for LiM2(PO4)3 and Li3M2(PO4)3 NASICON

compounds, where Li is located in one of the sites. On raising the temperature, Li

distribution within M1 cavity changes substantially, leaving the M1 site and moving

towards peripheral positions [61].

Ion transport in a solid occurs through the transfer of defects; ionic

conductivity is the product of defect concentration and mobility. Hence, heterovalent

doping increases the ion mobility in NASICON. Materials containing various

univalent cations show the minimum conductivity with composition independent of

temperature [62]. One of the high ion conducting NASICONs is prepared by partial

substitution of phosphorous by silicon in 1:2 ratio [63]. Compounds in which B, As,

Al or V partially substitutes for phosphorous are explored and optimization of the

conducting properties is achieved for Na1.8Ti2B1.2P1.8O12 [64]. Attempt is made to

explore the conductivity studies in Ti based materials, where phosphorous is replaced

by pentavalent elements like vanadium, antimony etc. [65]. Ti cation is substituted by

trivalent and pentavalent elements and the overall charge is less than five. In

Na1+xAlxTi2-x(PO4)3 and Na1+xEuxTi2-x(PO4)3 series, conductivity increases with

doping due to increase in charge carrier density [66].

17

The Na1-xNbxZr2-x(PO4)3 and Na1+2xMgxZr2-x(PO4)3 solid solutions are studied

by nuclear magnetic resonance spectroscopy and showed that Na atoms are in M1

positions. Partial substitution of Zr causes some of Na+ ion to move from these

positions to lower symmetry M2 positions [67]. The substitution of Mg or Zr for

chromium i.e., Na3+xCr2-xMgx(PO4)3 and Na3-xCr2-xZrx(PO4)3 decreases the

conductivity. The maximum conductivity is evidenced in NASICON phases

containing three sodium ions per formula unit [68].

For lithium systems, the disorder in cation sub-lattice increases the defect

concentration. The substitution of tetravalent ion for a trivalent ion generates lithium

site vacancies and conductivity increases by several orders of magnitude. Cation

mobility and disorder suppress the phase transition temperature. The phase transition

temperature is shifted by heterovalent substitution of Zr in LiZr2(PO4)3. 50% of low

temperature triclinic phase transforms into high temperature rhombohedral phase

when trivalent or pentavalent cations substitute for one percent of zirconium and 5% -

10% doping stabilizes the rhombohedral phase at room temperature [69]. The Li

NMR spectra and conductivity measurements at low temperatures signify an increase

in cation mobility due to generation of extra vacancies and insertion of extra lithium

interstitials at low dop ing levels. However, charge carrier concentration decreases at

higher pentavalent doping concentration and conductivity decreases [70]. With high

chromium concentration, intrinsic rhombohedral phase of LiHf2(PO4)3 is changed to

orthorhombic phase [71].

Lithium titanium phosphate crystallizes in rhombohedral symmetry and Ti

doping by heterovalent elements are investigated [72]. The compound crystallizes in

orthorhombic or monoclinic symmetry at higher doping concentrations. The

conductivity passes through a maximum with doping level and is reported for

Li1+xTi2-xAlx(PO4)3 (x=0.3) at 570K is 10-2 to 10-3Scm-1. The Li+ ions are disordered

between positions of M1 and M2 [73]. Moreover, total electrical conductivity strongly

depends on density of the compound and nature of grain-boundaries [74].

Nanostructured materials attracted great interest due to their enhanced

properties. Important property is the enhancement in ionic conductivity in

nanocrystalline form. In conventional polycrystalline solid electrolytes, grain-

boundaries partially block the ionic transport, an extra contribution to total

conductivity. But in nanostructured materials, grain-boundary diffusion is much

faster than grain diffusion and its volume fraction is more in nanomaterials [75].

18

Alternate synthesis techniques like wet chemical synthesis and sol-gel procedures are

explored for the synthesis of NASICON materials in nanocrystalline form [76].

Solution combustion synthesis is an effective low cost method for the production of

ceramic materials in nanocrystalline form, but this technique is rarely used for the

synthesis of NASICON materials [77].

1.6 Review of Olivine Based Lithium Metal Phosphates

1.6.1 Structure of Olivine

Olivine compounds exhibit high redox potential, fast Li+ ion transport,

excellent thermal stability and energy density comparable to that of conventional

lithium metal oxides. Moreover, anion substitution can alter the redox potential due to

two effects: (i) an inductive effect due to changes in the metal ion energy levels

resulting from the changes in the anion group and (ii) production of fewer or more

electrons shift the lithium concentration at which given redox reaction takes place.

Lithium orthophosphates, LiMPO4, where M=Ni, Fe, Mn and Co assume an

olivine related structure, consists of hexagonal closed packing of oxygen atoms with

Li+ and M2+

cations located in half of the octahedral sites and P5+

cations in 1/8 of the

tetrahedral sites [78]. The LiMPO4 olivine structure consists of PO4 tetrahedra with

M2+ ions on the corner sharing octahedral positions and Li+ ions on the edge sharing

octahedral positions, the latter running parallel to the b-axis. Figs. 1.2(a)-(b) show the

structure of LiNiPO4 in different orientations. Fig. 1.2(a) shows the chains (along the

c direction) of the corner sharing NiO6 octahedral layers. Fig. 1.2(b) shows edge

sharing LiO6 octahedral chains, which are cross linked by PO4 groups forming the

three dimensiona l network. In the hexagonal close packing array of oxide ions Li

resides in octahedral sites. Li+ ions on the edge sharing octahedral positions, it is

running parallel to the b-axis. These compounds generally crystallize in orthorhombic

(Pnma) symmetry, which accommodates four units of LiMPO4. The co-valence of M-

O bond differs with the degree of co-valence in the P-O bond. In the ordered olivine

structure, M3+/M2+ level lies lower than NASICON like frame-work, giving a higher

voltage versus lithium. Out of the olivine members, weak ferromagnetism is found for

LiNiPO4, LiMnPO4 and LiCoPO4 but not for LiFePO4 [79].

19

Figure 1.2: Olivine structure of LiNiPO4 showing Li ions (yellow spheres), PO4 tetrahedra (red) and NiO6 octahedra (blue) (a) Shows the chains (along the c-direction) of corner sharing NiO6 octahedral layers (b) Shows edge sharing LiO6 octahedral chains which are cross linked by PO4 groups forming the three dimensional network. (Courtesy: J. Gopalakrishnan et al.)

Atomistic simulation studies suggest that most favourable intrinsic defect in

bulk LiNiPO4 is the ‘anti-site’ defect, for which a small population (less than 2%) of

Li+ and Ni2+ ions are expected to exchange sites [80]. This defect would be

temperature dependent and sensitive to experimental synthesis conditions. Structural

analysis of hydrothermally synthesized LiFePO4 suggests 3mol% Fe on the lithium

sites, while scanning transmission electron microscopy study observed anti-site

defects in LiFePO4, quoting a concentration of around 1%. Nickel based olivine has

low capacities, less than 125mAhg-1; extremely low conductivities and high oxidation

potential of 5V make it unsuitable for use with known electrolytes. The potentials of

LiFePO4 and Li3Fe2(PO4)3 in lithium cells range from around 2.5 to 3.5V. This larger

difference between Fe2+/Fe3+ and Li/Li+ relative to iron oxides is explained by the

inductive effect [81]. The more electronegative the central atom in PO4 groups, higher

the cell potential. The presence of PO4 group in the structure causes flow of electron

density away from the M–O bond.

1.6.2 Conductivity in Olivine Compounds

In olivine form of lithium metal phosphates, octahedra are connected to form

the tunnel structure. These tunnels are along the b axis but are not connected, so the

20

lithium ions residing in the tunnels cannot readily jump from one tunnel to another.

Thus, lithium ion diffusion is one dimensional and any blockage in the tunnel

prevents movement of lithium ions [82]. Any attempt to dope these materials with

aliovalent ions must ensure that these ions do not reside in the lithium sites, but rather

in M sites. The exact mechanism of lithium ion motion in olivine type structures is

still not known, but an idea about the magnitude can be obtained from conductivity

measurements.

The electrical conductivity of pure LiFePO4, synthesized from the non-carbon

containing reactants, is less than 10-9Scm-1 [83]. Assuming the conductivity is solely

due to lithium ions and all the lithium ions are mobile. In pure LiFePO4, no

electronically conducting species is present, so that the total conductivity is as low as

10-11Scm-1[84] and the conductivity is due to motion of lithium ions. Many LiFePO4

samples are prepared using carbon containing precursors such as oxalates, acetates or

carbonates, which resulted in some residual carbon in the reducing atmospheres.

These samples typically have much higher conductivities, around 10-5 to 10-6Scm-1

but still too low for high power battery applications [85]. Addition of magnesium or

titanium resulted in poorer electrochemical performance, perhaps because of

occupancy of the lithium sites. Julien et al., reported the electrical conductivity and

micro structural properties of LiMPO4 system synthesized by solid-state reaction [86].

The strong covalent bonding between oxygen and P5+ to form (PO4)3- units

allow for greater stabilization of the structure compared to layered oxides, e.g.

LiCoO2, where, the oxide layers are more weakly bound. This strong co-valency

stabilizes the anti-bonding Fe3+/Fe2+ state through Fe-O-P inductive effect. Li

intercalation potential for LiNiPO4 is close to 5.1V, which is above the maximum

potential that can be tolerated by Li electrolytes and delithiation of LiNiPO4 is not

observed.

1.7 Computational Methods The occurrence of fast and inexpensive microcomputers during the past

decade has made computational techniques widely available and major progress is

seen in last years. The most common computational technique is the finite element

calculation: it has been applied to simulate the influence of various experimental

parameters on the impedance spectra of electroceramics, including nonhomogeneous

21

contacts and variation of grain-boundary properties. The dynamic and electronic

properties of mobile ions in the superionic solids are studied using molecular

dynamics simulation [87]. Numerical simulation techniques have been used to model

the conductivity of ionic conductor-insulator composites, especially near the

percolation threshold [88] and fractal decomposition structures are observed after

solid state electrolysis of CuBr films at large direct current voltage [89]. One major

goal of computational material science is the predictive capability. Ceder et al., [90]

gave an example to calculate the significant enhancement in the open circuit voltage

of LixCoO2 cathode materials by the addition of Al2O3. An overview on atomistic

computational material science based on quantum chemical concepts can be found in

an article by Hafner. Such techniques have been used recently to determine the

oxygen migration energies in CeO2 based oxides [91].

1.8 Ionic Conduction in the Bulk: Hopping Model

Ionic conductivity is due to the thermally activated ion hopping. D is the

diffusion coefficient, a is function of jump distance, ω0 is the characteristic attempt

frequency, Eh is the free energy of migration and the diffusion coefficient is given

by:

D=γa2ω0exp�- EhkBT� (1.2)

The factor γ takes into account geometrical and correlation effects. The backward

jump has higher probability than forward jump and cooperative motion may lead to

higher diffusion coefficients than isolated jumps. The mobility, µ, is related to

diffusion coefficient, D as:

µ= DqkBT

(1.3)

Using Eqs. (1.2) & (1.3), the ionic conductivity can be expressed as:

σdc=ncqµ=n0q2γa2ω0exp�- EfkBT� exp�- Eh

kBT� /kBT (1.4)

The creation of charge carriers is thermally activated and carrier density, nc, is of the

form:

nc=n0exp�- EfkBT� (1.5)

The general form of ionic conductivity, σdc, then becomes

σdc= σ0T

exp�− EσkBT� (1.6)

22

where, Εσ is the sum of the migration and formation energies i.e., Εσ=Εf+Εh. Most of

the crystalline and amorphous fast ionic conductors satisfy the Eq. (1.6), where Εf is

negligible.

1.9 Different Relaxation Functions for Understanding the ac Electrical Relaxation Process in Ion Conducting Materials

One of the important characteristic properties of electrical conduction in

disordered solids is dispersion in ac conductivity. In many non-metallic ionic

conductors dc or ac electrical conductivity is the result of diffusion of ions through the

conductors. As Almond et al., pointed out, very often the dispersive behavior of ac

conductivity of (disordered solids) ion conducting materials can be expressed in the

form [92]:

σ'(ω)=σdc+Aωn (1.7)

where, the first term σdc is the frequency independent dc conductivity and the second

term is purely dispersive component of the ac conductivity, depends on the

measurement frequency, ω, in a characteristic power law fashion. The frequency

exponent n lies in the range 0.5 to 1 and the parameter A is frequency independent,

but may be temperature dependent.

The concept of dipoles giving rise to polarization was introduced by Debye

and Debye response is obtained for an assembly of non-interacting ideal dipoles.

Solid electrolyte materials show a non-Debye dielectric response, characterized by a

broad loss peak and constant phase angle behaviour at frequencies above the loss peak

[93]. Hence, it is essential to have a proper theoretical understanding of rather general

behaviour.

In literature, relaxation functions like Cole-Cole, Cole-Davidson, Havriliak-

Negami, Kohlrausch-William-Watts functions are widely used for the analysis of

electrical relaxation process [94]. In general, non-Debye relaxation behaviour can be

regarded as arising from the elongation of time scale, i.e., a single effective relaxation

time is no longer sufficient to describe the entire relaxational behaviour. The various

empirical forms of different relaxation functions are [95]:

Cole-Cole function: Z*(ω)= 1[1+(𝑖ωτ)αcc]

(1.8)

Cole-Davidson function: Z*(ω)= 1

[1+(𝑖ωτ)]βCD (1.9)

23

Havriliak-Negamai function: Z*(ω)= 1[1+(𝑖ωτ)αcc]βCD (1.10)

Williams-Watts stretched exponential function: ϕww(t)= exp[-�𝑡𝜏�βww

] (1.11)

where, 0≤αcc<1, 0<βCD≤1 , these parameters accommodates the distribution of

relaxation times. The ω is angular frequency of applied electric field, τ is the mean

relaxation time. The position of dielectric loss peak is shifted to high frequency

(above Debye loss peak) with asymmetry in the loss peak curve in these relaxation

functions. In time domain relaxation function, the exponent, 0<βww≤1, ϕww(t) is

deviating from exponential decay for tτ≪1 and t

τ≫1 as βww decreases. The existing

non-Debye functions are either inadequate to explain the ac conductivity features or

fails to satisfy the linear response theory. Acquiring motivation from this inadequacy

of non-Debye functions, an anomalous relaxation function is proposed by Govindaraj

et al., [96]. The recently proposed model is capable of explaining the dispersion in

electrical relaxation, while satisfying the linear response theory criteria.

1.10 Present Work

With the outlook of available literatures, present study explored high energy

ball-milling and solution combustion technique for the synthesis of NASICON type

materials in nanocrystalline form. The olivine type materials are prepared in

nanocrystalline form by solution combustion technique. The following NASICON:

(i) LiTi2(PO4)3

Li1.3Al0.3Ti1.7(PO4)2.9(VO4)0.1

Li1.0Al0.4Ti1.7(PO4)2.88(VO4)0.12

Li1.6Al0.2Ti1.7(PO4)2.94(VO4)0.06

(ii) Na3Cr2(PO4)3: Glycine, urea and citric acid in 1:1, 1:2 and 1:3 molar ratios. Li3Fe2(PO4)3: Glycine in 1:2 molar ratio and citric acid-ethylene glycol in 1:1

molar ratio. Na3Fe2(PO4)3: Citric acid-ethylene glycol in 1:1 molar ratio. Li3Fe2-xCrx(PO4)3, where x=0.1, 0.2, 0.3, 0.4, 0.5: Citric acid-ethylene glycol in

1:1 molar ratio. Li3+xFe2-xMgx(PO4)3, where x=0.1, 0.2, 0.3, 0.4, 0.5: Citric acid-ethylene glycol

in 1:1 molar ratio.

24

and the following olivine: (iii) LiNiPO4: Citric acid in 1:2 molar ratio.

LiNi1-xCuxPO4, where x=0.1, 0.2, 0.3, 0.4, 0.5: Citric acid in 1:2 molar ratio.

LiNi1-xMgxPO4, where x=0.1,0.2, 0.3, 0.4, 0.5: Citric acid in 1:2 molar ratio.

systems are synthesised in the present study. In NASICONs, Li3Fe2-xCrx(PO4)3,

Li3+xFe2-xMgx(PO4)3 and in olivine, LiNi1-xCuxPO4, LiNi1-xMgxPO4 are new series and

are not reported in the literature.

LiTi2(PO4)3 is one of the well studied NASICON type material because of its

reasonably high ion conductivity at room temperature. In the present study, high

energy ball-milling technique is explored to increase the conductivity of

microcrystalline material by reducing crystallite size. The nanocrystalline material

shows one order increase in conductivity compared to the microcrystalline material,

due to feasible conduction through grain-boundaries [97, 98]. The average crystallite

size decreases with milling duration and the volume fraction of grain-boundaries

increases. The Al and V substituted LiTi2(PO4)3 material is ball-milled at different

durations and physical and electrical properties are investigated.

The Na3Cr2(PO4)3 and Li3Fe2(PO4)3 materials showed reversible structural

phase transitions in the microcrystalline form. The present study synthesizes these

materials in thermally stable nanocrystalline form by solution combustion technique.

Effects of nature of the fuel/chelating agent and its molar ratio on physical and

electrical properties of Na3Cr2(PO4)3 are investigated in detail. The electrical

properties are found to be dependent on physical properties of the material [99,100].

Nanocrystalline Li3Fe2(PO4)3 is synthesized using glycine and citric acid-ethylene

glycol mixture and is crystallized in orthorhombic crystal system with Pcan space

group. The improved dc conduction in Li3Fe2(PO4)3, synthesized using glycine, is

explained through enhanced hopping rate and unit cell volume. The dc conductivity of

Na3Fe2(PO4)3 is higher than Li3Fe2(PO4)3 due to rattling of ions within the large

interstitial space available. Li3Fe2(PO4)3 and Na3Fe2(PO4)3 are crystallized in

monoclinic symmetry i.e., β-Fe2(SO4)3 type structure. In this structure, mobile ions

occupy four co-ordination sites which are preferable to six co-ordination sites of

NASICON for ion conduction.

25

The effects of B-site ion radius and polarization on the conductivity are

investigated by substitution of Fe3+ in Li3Fe2(PO4)3 by Cr3+ and Mg2+ in different

compositions. The mobile ion density of Li3Fe2-xCrx(PO4)3 is almost same in the

series, but dc conductivity shows increasing trend with Cr3+ substitution, reaching a

maximum at x=0.3, owing to its high mobility. The Mg2+ substitution increases Li+

ion occupation rate and the carrier density is apparently the same for all compositions.

The substitution of Fe3+ by electropositive Mg2+ increases the ionic character of the

skeleton and as a consequence the Li-O bonds are strengthened; hence mobility

decrease with composition. The highest conductivity of x=0.1 composition is

attributed to its mobility.

LiNiPO4 based olivine materials are synthesized using citric acid in 1:2 molar

ratio. Solution combustion technique reduces the crystallite size to nanometer range

and the electrical conductivity of the sample is improved in nanocrystalline phase. To

investigate the effect of octahedral metal ion substitution; LiNi1-xMgxPO4 and

LiNi1-xCuxPO4 are (where x=0.1, 0.2, 0.3, 0.4 and 0.5) synthesized. In Mg and Cu

substituted series, highest conductivity is exhibited by x=0.1 composition. But

compared to the parent material, these compositions didn’t show any improvement in

conductivity. The decrease in conductivity in other compositions is due to the “anti-

site” defect, where a small population of Li+ and Ni2+ ions exchange their sites on the

substitution of aliovalent ions for Ni2+.

Even though, ac electrical data are analyzed using constant phase element,

power law and modulus formalism, the physical picture for the dispersion in ac

conductivity still needs further interpretation or an alternate approach. Recently

proposed alternate model [96] clearly discerns the fast process from slow relaxation

and extracts σdc, ωP and the exponent g. Slow and fast relaxation processes are related

to each other through the parameter g, which quantify the dispersion process in ion

conducting solids. Due to the trapping and correlated motions of charge carriers, fast

process is getting delayed. The extracted dc and ac electrical parameters of ion

conducting solids show good agreement with the experimental observation. The

model is successful in explaining the fast and slow relaxation process in glassy

systems. As a first attempt, this is demonstrated by analysing the NASICON glassy

systems like Li2.6+xTi1.4-xCd(PO4)3.4-x and K2.6+xTi1.4-xCd(PO4)3.4-x (where x=0.4, 0.8

and 1.0) [23].

26

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