prr.hec.gov.pkprr.hec.gov.pk/jspui/bitstream/123456789/2067/2/1603S.pdf · CONTENTS Acknowledgement...

Effect on Physical, Electrical and Magnetic Properties of

Magnesium Ferrite Nanomaterials Substituted by Different

Metal Cations

A dissertation submitted to the Department of Chemistry,

Quaid-i-Azam University, Islamabad, in partial fulfillment

of the requirement for the degree of

Doctor of Philosophy

In

Physical Chemistry

By

Zahoor Ahmad

DEPARTMENT OF CHEMISTRY

QUAID-I-AZAM UNIVERSITY

ISLAMABAD, PAKISTAN

2012

CONTENTS

Acknowledgement (i)

Abstract (ii)

List of units used (iv)

List of acronyms (vi)

Index of tables (viii)

Index of figures (x)

1 Introduction 01-43

1.1 Nanotechnology 01

1.1.1 Nanomaterials 02

1.2 Spinel compounds 04

1.2.1 Cation distribution in spinel compounds 05

1.3 Spinel ferrites 06

1.3.1 Chemical composition of spinel ferrites 06

1.3.2 Spinel magnesium ferrite 07

1.3.3 Crystal structure of spinel ferrites 08

1.4 Properties of magnesium ferrite 11

1.4.1 Magnetic properties 11

1.4.1.1 Magnetic parameters 12

1.4.1.2 Magnetic anisotropy 13

1.4.2 Electrical properties 14

1.4.2.1 DC-electrical properties 14

1.4.2.2 Dielectric properties 15

1.5 Advantages of magnesium ferrite 16

1.6 Applications of magnesium ferrite 17

1.7 Synthesis and characterization of magnesium ferrites: Literature 18

Review

1.8 Objectives and plan of work 42

2 Experimental 44-71

2.1 Methods of preparation 44

2.1.1 Microemulsion method 47

2.2 Chemicals used 48

2.2.1 Samples preparation procedure 48

2.3 Characterization techniques 50

2.3.1 Thermo-gravimetric analysis 51

2.3.2 Structural analysis 52

2.3.2.1 Powder X-ray diffraction (XRD) 53

2.3.2.2 Scanning electron microscopy (SEM) 55

2.3.2.3 Energy dispersive X-ray fluorescence (ED-XRF) 55

2.3.3 Mössbauer spectrometry 57

2.3.4 Magnetic properties measurement systems 60

2.3.4.1 SQUID magnetometer 60

2.3.4.2 Vibrating sample magnetometer (VSM) 63

2.3.5 Electrical and dielectric measuring systems 66

2.3.5.1 DC-electrical resistivity 66

2.3.5.2 Dielectric properties measuring system 69

3 Results and Discussion 72-149

3.1 Thermal properties 72

3.2 Structural properties 77

3.2.1 X-ray diffraction (XRD) analysis 77

3.2.2 Scanning electron microscopic (SEM) analysis 88

3.2.3 Energy dispersive X-ray fluorescence (ED-XRF) analysis 92

3.3 Mössbauer analysis 93

3.4 Magnetic measurements 104

3.4.1 SQUID magnetometric measurements 104

3.4.2 Vibrating sample magnetometric (VSM) measurements 122

3.5 Electrical properties 127

3.5.1 DC-electrical resistivity measurements 127

3.5.2 Dielectric measurements 138

Conclusions 147

Suggestions for further research 149

References 150-168

i

ACKNOWLEDGEMENT

I express my gratitude and appreciation to Prof. Dr. Muhammad Javed Iqbal for

supervision and for taking his keen interest and sincere counseling in every phase of

this research work. I would like to thank Prof. Dr. Saqib Ali, Chairman, Department

of Chemistry and Prof. Dr. Muhammad Siddiq, Head of Physical Chemistry, for

providing necessary research facilities.

Thanks are due to Higher Education Commission (HEC) of Pakistan for financial

support under Indigenous PhD 5000 Fellowship Scheme and International Research

Support Initiative Program (IRSIP). I am indebted to Dr. Turgut Meydan, Dr. Yevgen

Melikhov and Dr. Ikenna C. Nlebedim for hosting me as well as for sharing ideas and

opinions during my six month academic visit at Wolfson for Magnetics, School of

Engineering, Cardiff University Cardiff CF24 3AA, U.K.

This all is the fruit of lots of prayers and moral support of my mother and financial

support of my elder brother Haji Muhammad Musa. I cannot forget the lot of prayers

and sacrifices of my spouse and daughters Zainab and Areeba. In addition, I sincerely

acknowledge the lot of prayers and support of my sister, sisters in law and brothers in

law, nephews and nieces. I am heartily thankful to my parents in law for moral

support and especially for their patience to accommodate my family for long duration

of this higher study. Moral support of my teacher and Uncle Prof. Ghulam Rasool is

gratefully acknowledged.

I am thankful to all members of my research group particularly Qaisar, Naeem,

Rafaqat, Mahrukh, Mansoora, Saima, Bushra, Irfan and Waqas for their cooperation. I

am also thankful to my fellows Abbas, Arshad Khosa, Ansar Yaseen, Shahid Saeed,

Mahmood, Tashfeen, Usman, Afzal Shah, Rizwan and Iqbal for their affections and

nice companionship. Lastly, I would like to extend my heartfelt gratitudes to all my

friends particularly Nasr Ullah, Munawar, Shaukat Nadeem, Rehan, Prof. Hafeez

Ullah and Zahid whose concerns and encouragements really empowered me to

complete this work.

Zahoor Ahmad

ii

ABSTRACT

In the present study, M-Cr (M = Co, Ni, Cu, Zn and Mn) substituted magnesium

ferrite nanomaterials (Mg1-xMxCrxFe2-xO4 with x = 0.0-0.5) have been prepared by the

polyethylene glycol assisted micro-emulsion method. Thermal and XRD analyses

reveal that the complete spinel cubic phase formation occurred at 1123 K. The

average crystallite sizes in differently doped series are in the range of 15-62 nm. The

micrographs obtained from SEM analysis show that the synthesized materials are

agglomeration of the individual particles. The energy dispersive X-Ray fluorescence

(ED-XRF) spectrometric analysis reveals that the observed molar ratios of different

components of the samples are in close agreement with their nominal compositions.

Variation of Mössbauer parameters is explained on the basis of preferential site

occupancy of the substituted cations. The center shift (CS) value for A-site is smaller

than that of B-site due to difference in the Fe3+

-O2-

inter-nuclear separation, normally

larger for B-site as compared to that for A-site. The value of quadrupole splitting (QS)

is negligibly small which indicates that the overall symmetry of Fe3+

surroundings is

not disturbed with the substitution of the dopant ions into a magnesium ferrite matrix.

With the increase of dopant contents, the variations of hyperfine magnetic field (H)

and site population area (A) are akin to the compositional variation of saturation

magnetization, MS. Symmetric magnetic hysteresis loops are measured using a

superconducting quantum interference device (SQUID) magnetometer up to an

applied magnetic field of 50 kOe at 300, 200 and 100 K. SQUID analysis reveals

narrow hysteresis loops with a coercivity (HC) and saturation magnetization (MS)

varying for different compositions. The high field regimes of these loops are modeled

using the Law of Approach (LoA) to saturation to extract information regarding

magnetocrystalline anisotropy and saturation magnetization. In the present study, the

saturation magnetization of magnesium ferrite increases by doping with Co-Cr, Ni-Cr,

Zn-Cr and Mn-Cr, respectively, but decreases by doping with Cu-Cr contents. The

coercivity (HC) of all the series studied here decreases with an increase in the

substitution level. All the magnetic parameters i.e. MS, Mr, K1 and HC increase with

decrease in the temperature from 300 K to 100 K. To determine the Curie temperature

(TC), the temperature dependence of normalized moment is measured at an applied

field of 5 kOe within the temperature range of 350-973 K, using a vibrating-sample

iii

magnetometer. The Curie temperature initially increases with Co-Cr and Mn-Cr

contents, but start to decrease for higher level of substitution. In case of Ni-Cr doped

series, TC value increases progressively with the increase in dopant contents, while it

continued to decrease with the substitution of Cu-Cr and Zn-Cr contents. Temperature

dependence of DC-electrical resistivity reflects the semi-conducting nature of the

doped Mg-ferrites. The room temperature resistivity (ρRT

) and activation energy

increase up to a certain level of substitution with Ni-Cr, Zn-Cr and Mn-Cr contents

but increase continuously for Co-Cr and Cu-Cr substitution contents. The dielectric

constant (έ) and dielectric loss tangent (tan δ) decrease with increasing applied field

frequency. The variations in the magnitude of drift mobility, dielectric constant and

dielectric loss tangent are in close agreement with the trend of DC-electrical

resistivity by increasing the dopant contents. With improvement in properties, the

synthesized materials could be suitable for potential application in some magnetic and

microwave devices.

iv

LIST OF UNITS USED

Å: Angstrom

Å3: Angstrom cube

ºC: Degree celsius

cm2V

-1s

-1: Centimeter square/volt second

emu: Electromagnetic unit

emu/g: Electromagnetic unit/gram

eV: Electron volt

g/cm3: Gram/centimeter cube

h: Hour

Hz: Hertz

J/m3: Joule/meter cube

JK-1

mol-1

: Joule/Kelvin. Mole

K: Kelvin

kA/m: Kilo ampere/meter

keV: Kilo electron volt

kHz: Kilo Hertz

kJ/mol: Kilo Joule/mole

kOe: Kilo oersted

m: Meter

MA/m: Mega ampere/meter

MeV: Mega electron volt

mgK-1

: Milligram/kelvin

MHz: Mega Hertz

min: Minute

mm/s: Millimeter/second

molL-1

: Mole/liter

ms: Milli second

µm: Micrometer

µs: Micro second

µV: Micro volt

nm: Nanometer

v

Oe: Oersted

Ω.cm: Ohm centimeter

pm: Pico meter

ppm: Parts per million

R: Gas constant

T: Tesla

V: Volt

W: Watt

vi

LIST OF ACRONYMS

AAS: Atomic absorption spectroscopy

AC: Alternating current

BSE: Backscattered electrons

CMC: Critical micelle concentration

CRT: Cathode-ray tube

CTAC: Cetyltrimethyl ammonium chloride

DC: Direct current

DSC: Differential scanning calorimetry

DTG: Differential thermal gravimetry

DTGS: Deuterium tryglycine sulfate

ED-XRF: Energy dispersive X-ray fluorescence

EPR: Electron paramagnetic resonance

ESR: Electron spins resonance

FESEM-EDS: Field emission scanning electron microscopy-energy dispersive

spectroscopy

FTIR: Fourier transforms infrared spectroscopy

HR-TEM: High resolution transmission electron microscopy

HT-XRD: High temperature X-ray diffraction

ICSD: International crystal system database

IR: Infrared

JCPDS: Joint committee on powder diffraction standards

LCR: Inductance Capacitance Resistance

ln: Natural logarithm

LoA: Law of approach

MPMS: Magnetic property measurement system

MW: Microwave

M(OR)x: Metal alkoxides

PEG: Polyethylene glycol

PVP: Polyvinyl pyrrolidone

R = Resistance

RH: Relative humidity

vii

SDS: Sodium dodecyl sulfate

SE: Secondary electrons

SEAD: Selected-area diffraction

SEM: Scanning electron microscopy

SQUID: Super conducting quantum interference device

TC: Curie temperature

TEM: Transmission electron microscopy

TGA: Thermo-gravimetric analysis

UV: Ultraviolet

UV/VIS: Ultraviolet/ visible spectroscopy

VSM: Vibrating sample magnetometer

XRD: X-ray diffraction

viii

INDEX OF TABLES

Table 2.1 Specifications for the chemicals used 48

Table 3.1 Crystallite size (D), lattice constant (a), cell volume (Vcell), X-ray

density (dx) and bulk density (db) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5)

80


density (dx) and bulk density (db) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5).

86


density (dx) and bulk density (db) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5).

86


density (dx) and bulk density (db) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5).

87


density (dx) and bulk density (db) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5).

87

Table 3.6 The observed and theoretical composition of selective samples of

Mg1-xMxCrxFe2-xO4 (M = Co, Ni, Cu, Zn, Mn, and x = 0.0, 0.2 and 0.4).

92

Table 3.7 Center shift (CS), quadrupole splitting (QS), hyperfine magnetic field

(H) and relative area (A) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5).

95


(H) and relative area (A) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5).

102


(H) and relative area (A) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5).

102


(H) and relative area (A) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5).

103


(H) and relative area (A) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5).

103

Table 3.12 Saturation magnetization (MS), remanence (Mr), magnetocrystalline

anisotropy constant (K1) and coercivity (HC) of Mg1-xCoxCrxFe2-xO4

(x = 0.0-0.5) at 300, 200 and 100 K.

106


anisotropy constant (K1) and coercivity (HC) of Mg1-xNixCrxFe2-xO4

(x = 0.0-0.5) at 300, 200 and 100 K.

118


anisotropy constant (K1) and coercivity (HC) of Mg1-xCuxCrxFe2-xO4

(x = 0.0-0.5) at 300, 200 and 100 K.

119

ix

Table 3.15 Saturation magnetization (MS), remanence ( Mr), magnetocrystalline

anisotropy constant (K1) and coercivity (HC) of Mg1-xZnxCrxFe2-xO4

(x = 0.0-0.5) at 300, 200 and 100 K.

120


anisotropy constant (K1) and coercivity (HC) of Mg1-xMnxCrxFe2-xO4

(x = 0.0-0.5) at 300, 200 and 100 K.

121

Table 3.17 Curie temperature (TC) of Mg1-xMxCrxFe2-xO4 (M = Co, Ni, Cu, Zn, Mn

and x = 0.0-0.5)

124

Table 3.18 DC-electrical resistivity (ρRT

), activation energy (Ea), drift mobility

(µd), dielectric constant (έ) and dielectric loss tangent (tanδ) of

Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5).

130




Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5).

136




Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5).

137



(µd), dielectric constant (έ) and dielectric loss tangent (tan δ) of

Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5).

137




Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5).

138

x

INDEX OF FIGURES

Figure 1.1 Schematic of partial unit cell 8

Figure 1.2 Schematic drawings of lattice surroundings and nearest neighbors for

(a) the tetrahedral A-site (8a), (b) the octahedral B-site (16d), and (c)

the tetrahedral oxide site (32e). Anion dilations are indicated in (a)

by solid arrows

10

Figure 1.3

Figure 1.4

Different types of magnetism in spinel ferrites

The hysteresis loop of a magnetic material, where H is the magnetic

field amplitude and M is the magnetization of the material

11

13

Figure 1.5 Electric field interactions with an atom under the classical dielectric

model

16

Figure 2.1 Flow sheet diagram for the synthetic scheme 49

Figure 2.2 A block diagram of a thermo-gravimetric analyzer 52

Figure 2.3 Illustration of crystal planes and Bragg’s law 54

Figure 2.4 Block diagram for energy dispersive X-ray fluorescence 56

Figure 2.5 The Chemical Shift (CS) and Quadrupole Splitting (QS) of nuclear

energy levels and the corresponding Mössbauer spectrum

58

Figure 2.6 A schematic view of typical Mössbauer spectrometer 59

Figure 2.7 Basic scheme of MPMS Quantum Design 60

Figure 2.8 SQUID detection diagram 61

Figure 2.9 A symmetric hysteresis loop of magnesium ferrite measured by the

SQUID magnetometer

63

Figure 2.10 Flow sheet for the working of vibrating sample magnetometer

(VSM)

64

Figure 2.11 Moment signals to find out the saddle point of the sample 65

Figure 2.12 Demagnetization curves of nickel sample as a function of

temperature at two temperature sweep speeds

65

Figure 2.13 Flow sheet diagram of the two-probe resistivity apparatus 67

Figure 2.14 Temperature dependence of electrical resistivity measured by the

device described in Sec. 2.3.5.1

68

Figure 2.15 Arrhenius type relationship for activation energy calculation 69

Figure 2.16 Schematic diagram of LCR meter 70

xi

Figure 2.17 Frequency dependence of the dielectric constant measured by the

apparatus described in Sec. 2.3.5.2.

71

Figure 3.1 TGA/DTG curves of the as-synthesized pure magnesium ferrite 73

Figure 3.2 TGA/DTG curves of the as-synthesized Mg0.6Co0.4Cr0.4Fe1.6O4 74

Figure 3.3 TGA/DTG curves of the as-synthesized Mg0.6Ni0.4Cr0.4Fe1.6O4 75

Figure 3.4 TGA/DTG curves of the as-synthesized Mg0.6Cu0.4Cr0.4Fe1.6O4 75

Figure 3.5 TGA/DTG curves of the as-synthesized Mg0.6Zn0.4Cr0.4Fe1.6O4 76

Figure 3.6 TGA/DTG curves of the as-synthesized Mg0.6Mn0.4Cr0.4Fe1.6O4 76

Figure 3.7 XRD patterns of the synthesized MgFe2O4 sample compared with the

standard pattern

77

Figure 3.8 XRD patterns of Mg-ferrite doped with Cox-Crx contents (x = 0.0-

0.5)

79

Figure 3.9 XRD patterns of Mg-ferrite doped with Nix-Crx contents (x = 0.0-

0.5)

82

Figure 3.10 XRD patterns of Mg-ferrite doped with Cux-Crx contents (x = 0.0-

0.5)

83

Figure 3.11 XRD patterns of Mg-ferrite doped with Znx-Crx contents (x = 0.0-

0.5)

84

Figure 3.12 XRD patterns of Mg-ferrite doped with Mnx-Crx contents (x = 0.0-

0.5)

85

Figure 3.13 Scanning electron micrograph of MgFe2O4 sample 88

Figure 3.14 Scanning electron micrographs of Mg1-xCoxCrxFe2-xO4 (x = 0.1, 0.3) 89

Figure 3.15 Scanning electron micrographs of Mg1-xNixCrxFe2-xO4 (x = 0.1, 0.3) 90

Figure 3.16 Scanning electron micrographs of Mg1-xCuxCrxFe2-xO4 (x = 0.1, 0.3) 90

Figure 3.17 Scanning electron micrographs of Mg1-xZnxCrxFe2-xO4 (x = 0.1, 0.3) 91

Figure 3.18 Scanning electron micrographs of Mg1-xMnxCrxFe2-xO4 (x = 0.1, 0.3) 91

Figure 3.19 Mössbauer spectra of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) 94

Figure 3.20 Mössbauer spectra of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5) 98

Figure 3.21 Mössbauer spectra of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5) 99

Figure 3.22 Mössbauer spectra of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) 100

Figure 3.23 Mössbauer spectra of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5) 101

Figure 3.24 M-H loops of Mg0.8Co0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset:

First quadrant of magnetic hysteresis loops for Mg1-xCoxCrxFe2-xO4

105

xii

at 300 K

Figure 3.25 M-H loops of Mg0.8Ni0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset:

First quadrant of magnetic hysteresis loops for Mg1-xNixCrxFe2-xO4 at

300 K

115

Figure 3.26 M-H loops of Mg0.8Cu0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset:

First quadrant of magnetic hysteresis loops for Mg1-xCuxCrxFe2-xO4

at 300 K

116

Figure 3.27 M-H loops of Mg0.8Zn0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset:

First quadrant of magnetic hysteresis loops for Mg1-xZnxCrxFe2-xO4

at 300 K

116

Figure 3.28 M-H loops of Mg0.8Mn0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset:

First quadrant of magnetic hysteresis loops for Mg1-xMnxCrxFe2-xO4

at 300 K

117

Figure 3.29 Thermal variation of normalized moment in Mg1-xCoxCrxFe2-xO4

(x = 0.0-0.5)

122

Figure 3.30 Thermal variation of normalized moment in Mg1-xNixCrxFe2-xO4

(x = 0.0-0.5)

125

Figure 3.31 Thermal variation of normalized moment in Mg1-xCuxCrxFe2-xO4

(x = 0.0-0.5)

125

Figure 3.32 Thermal variation of normalized moment in Mg1-xZnxCrxFe2-xO4

(x = 0.0-0.5)

126

Figure 3.33 Thermal variation of normalized moment in Mg1-xMnxCrxFe2-xO4

(x = 0.0-0.5)

126

Figure 3.34 Plot of DC-electrical resistivity of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5)

versus temperature; Inset: Arrhenius type plot of ln ρ vs. 103/T for

some Co-Cr substituted magnesium ferrite samples

128

Figure 3.35 Plot of DC-electrical resistivity of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5)


some Ni-Cr substituted magnesium ferrite samples

134

Figure 3.36 Plot of DC-electrical resistivity of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5)


some Cu-Cr substituted magnesium ferrite samples

135

Figure 3.37 Plot of DC-electrical resistivity of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) 135

xiii


some Zn-Cr substituted magnesium ferrite samples

Figure 3.38 Plot of DC-electrical resistivity of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5)


some Mn-Cr substituted magnesium ferrite samples

136

Figure 3.39 Dielectric constant (έ) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f

139

Figure 3.40 Dielectric loss tangent (tanδ) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5)

versus ln f

139

Figure 3.41 Dielectric constant (έ) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5) versus ln f 142

Figure 3.42 Dielectric loss tangent (tanδ) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5)

versus ln f

143

Figure 3.43 Dielectric constant (έ) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f

143

Figure 3.44 Dielectric loss tangent (tanδ) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5)

versus ln f

144

Figure 3.45 Dielectric constant (έ) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f

144

Figure 3.46 Dielectric loss tangent (tanδ) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5)

versus ln f

145

Figure 3.47 Dielectric constant (έ) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f

145

Figure 3.48 Dielectric loss tangent (tanδ) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5)

versus ln f

146

CHAPTER 1

INTRODUCTION

1

1. INTRODUCTION

1.1 Nanotechnology

Nanotechnology controls the structure of matter at the nanoscale to produce new

materials and devices with unique properties. However, some of these technologies have

limited control over structure at the nanoscale, but these are being used to produce useful

products. These are also being further developed to produce more sophisticated products

with structure in controlled manners. Nanotechnology is very diverse field, ranging from

extensions of conventional device physics to completely new approaches based upon

molecular self-assembly and development of new materials with nanoscale dimensions to

investigate the direct control on the atomic level. Nanotechnology involves application of

several scientific fields comprises biomedical sciences, surface science, electronics,

semiconductor physics, optics, magnetism, energy storage and electrochemistry [1].

There are many debates on the future implications of nanotechnology.

Nanotechnology may be able to produce many novel materials and devices with a vast

range of applications in medicine, biomaterials, electronics and energy production. On

the contrary, nanotechnology raises many questions like any advance technology

regarding toxicity and environmental impact of nanomaterials [2], their potential effects

on the global economy, and speculation about the several doomsday scenarios.

Although, nanotechnology is recent advances in scientific research, however the

development of its main concepts existed for a long period of time. The nanotechnology

emergence in 1980s was due to convergence of advanced experiments like the invention

of the scanning tunneling microscope in 1981 and fullerenes in 1985.

A nanometer (nm) is one billionth, or 10-9

of a meter. In comparison, the typical

length of carbon-carbon bond is in the range of 0.12 to 0.15 nm and the DNA double

helix has a diameter of about 2 nm. By convention, nanotechnology is considered in the

scale range of 1 to 100 nm according to the definition used by the National

Nanotechnology Initiative United States. The lower limit is determined by the size of

atoms (hydrogen has the smallest atoms, which are approximately one quarter of nm in

diameter). The upper limit is more or less arbitrary, but is approximately the size of the

phenomena not observed in larger structures start to become apparent and can be used for

nano devices [3].

2

Nanotechnology not only created many high quality products at very low cost, but

also allows producing nanofactories in the same low cost and at very fast speed.

At the nanoscale, electrical and magnetic properties are not the same as those of

their bulk counterparts. For example, opaque substances turn into transparent (copper);

stable materials turn combustible (aluminum); insoluble materials become soluble (gold).

In general a chemically inert material such as gold can be tuned as a potent chemical

catalyst at nanoscale.

There are many ways used to create nanostructured materials, which are usually

divided into two main strategies, top-down approach and bottom-up approach. The

traditional top-down approach is limited by the miniaturization and precise

manufacturing of semiconductor products at a smaller scale. Alternatively, in bottom-up

approaches, nanostructured materials are created from building blocks of atoms,

molecules, clusters or nanoparticles in a controlled manner, governed by thermodynamic

methods or new concepts such as self-assembly. Therefore, the idea of creating artificial

substances and materials with unique features by using the bottom-up approach is

increasingly encouraged for the development of new and multi-functional materials.

1.1.1 Nanomaterials

Over the past two decades, a class of materials with a nanosized microstructure

were prepared and studied. Nanomaterials have grain sizes on the order of a billionth of a

meter. All materials are composed of grains, which in turn comprised many atoms.

These grains are usually invisible to the naked eye, depending on their size. Conventional

materials have grains varying in size from 100’s of microns (µm) to millimeters (mm).

An average human hair is about 100 microns in diameter. The average size of an atom is

on the order of 1 to 2 angstroms (Å) in radius. 1 nanometer comprises 10 Å, and hence in

one nm, there may be 3-5 atoms, depending on the atomic radii. The nanomaterials are

assembled from nanoscale building blocks, mostly crystallites. The building blocks can

be different in their atomic structure, chemical composition and crystallographic

orientation. In cases, where the building blocks are crystallites, incoherent or coherent

interfaces can be formed between them, depending on the atomic structure, the

crystallographic orientation and the chemical composition of adjacent crystallites.

3

Recently, the synthesis of nanoscale magnetic materials has been an area of

intense study, due to new mesoscopic properties shown by quantum-sized particles

located in the transition region between atoms and bulk solids [4]. Quantum size effects

and the large surface area of magnetic nanoparticles dramatically change some of the

magnetic properties and present superparamagnetic phenomena and quantum tunneling of

magnetization. Several research groups are engaged in investigations of the metal oxide

nanoparticles because of their technological applications in magnetic and microwave

devices, magnetic recording media, etc. Several types of nanomaterials such as metal (Fe,

Co, Ni), metallic alloys (Fe-Cu) and metallic oxides (MgFe2O4, CoFe2O4, MnFe2O4 and

ZnFe2O4) are under recent research activity, while metal nanoparticles have stability

problems in atmospheric conditions. However, metal oxides are stable under ambient

conditions.

Nanocrystalline materials are exceptionally strong, hard, and ductile at high

temperature, wear-resistant, erosion-resistant, corrosion-resistant, and chemically very

active. Nanomaterials are also much more malleable than their conventional counterparts

commercially available. The deviation of properties of the nano sized materials from

those of bulk material is due to surface effects that depend primarily on the ratio of

surface area to volume and particle size, together with the chemical composition and

interaction between the particles.

The nanomaterials are classified into seven main categories [5].

1. Carbon based nanomaterial

2. Nanocomposites

3. Metals & alloys

4. Biological nanomaterial

5. Nano-polymers

6. Nano-glasses

7. Nano-ceramics

4

1.2 Spinel Compounds

Spinel phase crystallizes into the cubic system with octahedral crystal formation.

There are at least 30 oxides of minerals which included in the spinel super-group. The

majority of spinel compounds has space group mFd3 . The primary member of the group

has the general formula, AB2O4; wherein “A” represents a divalent metal ion such as

magnesium, manganese, nickel and zinc, etc. The “B” represents trivalent metal ions

such as aluminum, iron, chromium, etc. However, a binary mixture of titanium Ti+4

and

Pb+2 etc., can also occupy the octahedral site. Solid solution is quite common in this

group of minerals which means that it might contain specific percentages of various ions

in any particular specimen [6]. In most of oxide structures, the oxygen ions are

significantly larger than the cations having a spinel structure which can be assumed by a

cubic close packing of O2-

ions in which the cations (e.g. Mg2+

, Fe3+

) are located at

certain interstices. The spinel structure is similar to that of highly symmetric diamond

structure. The position of the “A” ions is almost identical to those occupied by the carbon

atoms in the diamond structure. This could explain the relatively good hardness and high

density of this typical group.

More than one hundred compounds of spinel structure reported to date. Most of

them are oxides, some are sulphides, selenides and telluride and few are halides. A large

variety of cations might be introduced into the spinel structure and different charge

combinations are possible; almost any combination that has eight positive charges to

balance eight anionic charges [7], for example, Mg2+Fe23+O4, Mg2+Al2

3+O4, Mg22+Ti4+O4,

Li1+

Al3+

Ti4+

O4, Li0.51+

Al2.53+

O4, and Na21+

W6+

O4 etc. In spinel oxide, normally the

different cations do not have a big difference in size, because the spinel structure is stable

only if the cations have rather medium size and in addition, the ionic radii of the different

metal species in the same compound are comparable and do not differ too much. Similar

cation combinations are presented in sulphides e.g. Zn2+Al23+S4 and Cu2

2+Sn4+S4.

However, in such spinels halide e.g. Li21+

Ni3+

F4 and Li1+

Mn23+/ 4+

O4, cations are limited

to valance state of +1 and +2, in order to exhibit an overall cation: anion ratio of 3: 4.

5

1.2.1 Cation distribution in spinel compounds

Spinels are classified on the basis of the position of cations in the two principal

sites, tetrahedral site (A) and octahedral site (B) [8], into three types, as described below:

i) Normal spinel

Normal spinel, M2+

(M23+

) O4, has all the divalent cations on the tetrahedral (A-)

sites and the trivalent cations on the octahedral (B-) sites. This can be represented by the

formula (M2+

)tet [M23+

]oct O42-

. Some examples are as follows.

CoO.Al2O3 = CoAl2O4 (normal)

ZnO.Fe2O3 = ZnFe2O4 (normal)

FeO.Al2O3 = FeAl2O4 (normal)

ii) Inverse spinel

The ‘inverse’ spinel, M3+

(M2+

M3+

) O4, has the divalent cations occupying the B-

sites and the trivalent cations are equally divided between A- and remaining B-sites. This

can be represented by the formula, (M3+

)tet [M2+

M3+

]oct O42-

. Examples;

CoO.Fe2O3 = CoFe2O4 (inverse)

NiO.Fe2O3 = NiFe2O4 (inverse)

iii) Intermediate spinel

In addition to the two extremes, i.e. normal and inverse spinels, intermediate

cation distribution is possible, represented as (M1-δ2+Mδ

3+)[Mδ2+M2-δ

3+]O42-

. The cation

distribution has been quantified by using a parameter δ, which corresponds to the fraction

of M3+

ions in the A- and B-sites. The small brackets represent the average occupancy of

A-sites, whereas the square brackets represent the average occupancy of B-sites. The

variable δ is the inversion parameter, which specifies the fraction of A-sites occupied by

M3+ ions [9].

Normal (M2+

)tet [M3+

M3+

]oct O42-

δ = 0

Inverse (M3+

)tet [M2+

M3+

]oct O42-

δ = 1

Intermediate M1−δ2+

Mδ3+

[Mδ2+

M2−δ3+

]O42- 0< δ<1

6

The inversion parameter is a measure of the degree of inversion and in some

ferrites depends on the method of preparation [10]. The factors affecting the cation

distribution over A- and B-sites are as follows [11, 12]:

• Ionic size of cations

• Electromagnetic configuration of the cations

• Electronic energy

Smaller cations prefer to occupy the A-sites. The cations have special preference for A-

and B-sites and the preference depends upon the following factors:

• Ionic radius

• Size of interstices

• Temperature

• Orbital preference for the specific coordination

The preference of cations is according to Verwey-Heilman scheme [13, 14]:

• Ions having a preference for A-sites: Zn2+

, Cd2+

, Ga2+

, In3+

, Ge4+

• Ions having a preference for B-sites: Ni2+

, Cr3+

, Ti4+

, Sn4+

• Ions having no specific site preference: Mg2+, Mn2+, Cu2+, Fe2+, Fe3+, Al3+

Moreover the electrostatic energy also affects the cation distribution in the spinel

lattice. The cations of the smallest positive charge reside on the B-sites having six anions

in surrounding i.e. the most favorable electrostatic conduction.

1.3 Spinel Ferrites

1.3.1 Chemical composition of spinel ferrites

Metal oxides with spinel structure often called “spinel” belong to the group of

strategic materials that are widely used in the modern technologies. They have excellent

magnetic, semiconducting, refractory, catalytic and sorption properties.

The general chemical formula of ferrites having the structure of the mineral spinel

(MgAl2O4) is MFe2O4, where as M represented a divalent metal ion with an ionic radius

of approximately 0.6-1.0 Å. In the case of simple spinel ferrites, M can be any of

transition metal including Mn, Co, Ni, Cu and Zn, or Mg and Cd. A combination of these

7

ions is also possible, called mixed ferrites. The symbol M might be a combination of ions

that has an average valency of +2 e.g. Li1+

and Fe3+

in lithium ferrite, Li0.5Fe2.5O4.

The trivalent iron ions (Fe3+) in MFe2O4 can be completely or partially substituted

by other trivalent metal ion like Al3+

or Cr3+

, resulting in a mixed crystals with

aluminates or chromites. These compounds can also behave as a ferrimagnetic at room

temperature provided that large amount of non-magnetic ions should not be incorporated.

If iron ions are replaced by a tetravalent ion like Ti4+

, an equal share of Fe3+

are

converted into Fe2+ to maintain the overall electro-neutrality. A wide variety of various

chemical composition of ferrimagnetic oxide with spinel structure is possible.

1.3.2 Spinel magnesium ferrite

Magnesium ferrite belongs to a class of soft ferrites having the general formula

MFe2O4 with spinel cubic structure. Magnesium ferrite (MgFe2O4) is a partially inverted

spinel ferrite, i.e. (Mg1-λFeλ)[MgλFe2-λ]O4 where parenthesis and square brackets denote

cation sites of tetrahedral (A-sites) and octahedral [B-sites] coordination, respectively

[15]. Mg2+

is a non-magnetic ion and has no contribution in the magnetic moment of

MgFe2O4, which is thus entirely due to the uncompensated spins of the un-evenly

distributed iron ions at two (A & B) sites. Magnesium ferrite is a typical spinel in which

the cation distribution in the crystal lattice site is very much sensitive to heat treatment

due to high diffusibility of Mg2+

ions [16]. The interesting physical and chemical

properties of ferrospinels arise from their ability to distribute the cations among the

tetrahedral (A) and octahedral (B) sites [17]. Further, various studies have shown that

magnetic and electrical properties are strongly linked to the structural properties, which

are controlled by the synthesis method. It also shows that substitution of different metal

cations has different effects on the distribution of cations within the spinel lattice. As a

result, a cation adjustment in spinel ferrites is a suitable approach for tuning magnetic and

electrical properties to suit intended applications. The change in magnetization and

conduction due to change in composition depends on the site occupancy and magnetic

moment contribution and electron hopping of the cations, its analysis can be used to

deduce trends in site occupancy with substitution.

8

1.3.3 Crystal structure of spinel ferrites

The crystal structure of spinel ferrite, MFe2O4 (M refers to the metal) can be

described as a cubic close-packed arrangement of oxygen atoms, with M2+

and Fe3+

distributed over two different types of crystallographic sites. These sites have tetrahedral

and octahedral coordination with oxygen (termed as A and B-sites, respectively), so that

the resulting local symmetry of both sites is different (Fig. 1.1).

Figure 1.1: Schematic of partial unit cell [18]

Spinel structure composed of two types of sites for cation occupation. These are

known as tetrahedral (A) and octahedral (B) sites. There are 8 A-sites wherein metal ions

have tetrahedral coordination with oxygen, and on tetrahedral site, the interstitial is at the

center of a tetrahedron formed by the four lattice atoms. Three adjacent atoms are in a

plane; the fourth atom is located at the top symmetrical position. In addition, the

tetrahedral site with defined geometry provides a space for an interstitial atom. On the

other hand, 16 B-sites which exhibits octahedral coordination offers a position to an

interstitial atom at the space in the interstices between 6 atoms forming regular

octahedron. Four regular atoms are positioned in a single plane; the remaining two are

located at symmetrical positions just above or below. All formed spheres can be

considered hard and adjacent to each other. The six spheres define a regular octahedron,

9

inside there is a space defined for an interstitial atom, surrounded by six spheres.

Consequently, there are 8 formula units per cubic unit cell (Z = 8), each consist of 32

anions and 24 cations, for a total of 56 atoms [9]. There are 96 interstices between the

anions in the cubic lattice; but in spinel ferrites, only 24 are occupied by cations. Out of

the 64 tetrahedral interstices (8a, 8b, 48f) between the anions, only 8 are occupied by the

cations. The rest of 16 cations occupied half of the 32 octahedral interstices (16c, 16d).

The unoccupied sites are octahedral (16c) and tetrahedral (8a, 48f) [20] as shown in Fig.

1.2.

The tetrahedral and octahedral sites have always fixed position and do not depend

upon the nature of the constituent cations. However, the general position of anions

depends on the relative size of A and B cations. The anion sub-lattice is arranged in a

pseudo-cubic close-packed (ccp) spatial arrangement, although some spinel possesses

almost-ideal ccp anion sub lattices. The repeat unit of the conventional unit cell is twice

that of the anion lattice. Accordingly, the spinel lattice parameter ‘a’ is large.

10

Figure 1.2: Schematic drawings of lattice surroundings and nearest neighbors for (a) the

tetrahedral A-site (8a), (b) the octahedral B-site (16d), and (c) the tetrahedral oxide site

(32e). Anion dilations are indicated in (a) by solid arrows [19].

48f tetrahedral vacancy A-site cation Oxide anion

2nd n.n. oxide ion

2nd n.n. B cation B-site cation 16c octahedral vacancy

8b tetrahedral vacancy

11

1.4 Properties of Magnesium Ferrite

The Mg-ferrites have attracted the attention of researchers since past few decades

[21-24]. One reason for this is the great potential of this material for applications, which

are based on higher values of saturation magnetization, Curie temperature and electrical

resistivity together with low dielectric losses and moderate coercive field. The ordering

of the magnetic moments of ferric ions and the strong exchange interactions explain the

excellent magnetic behavior of this material [25].

1.4.1 Magnetic properties

The magnetism of the spinel ferrite is due to uncompensated electron spins of the

individual magnetic ions and anti-parallel spin alignment in the two (A and B) sub

lattices of spinel structure [26].

Magnetic materials are classified by their response under the external applied

magnetic field. Descriptions of arrangement of the magnetic moments are helpful to

differentiate the different forms of magnetism found in materials (Fig. 1.3).

Figure 1.3: Different types of magnetism in spinel ferrites [18].

12

Five basic types of magnetism can be described: diamagnetism, paramagnetism,

ferromagnetism, antiferromagnetism and ferrimagnetism. In the presence of an externally

applied magnetic field the atomic current loops are created by the spinning of electrons

respond to repel the applied field. All materials possess such type of weak repulsion to an

external applied magnetic field called diamagnetism. All other types of magnetism

observed in the materials are almost due to unpaired electrons often in the 3d or 4f shells

of each atom. Materials with uncoupled magnetic moments exhibit paramagnetism.

Materials with aligned magnetic moments of equal magnitude possess ferromagnetism

and their crystalline structure offers the direct coupling interactions between the

moments, leading to an improvement in the flux density (e.g. Fe, Ni and Co). Materials

having atomic magnetic moments of equal magnitude oriented in an antiparallel way

possesses antiferromagnetism (e.g., troilite FeS and ilmenite FeTiO2).

The size reduction in magnetic materials results in the formation of single-domain

particles and materials show super-paramagnetism behavior. Each particle exhibits

paramagnetic nature with a giant magnetic moment, since there is still a well-defined

magnetic order in each of nanoparticles [27].

1.4.1.1 Magnetic parameters

The most commonly measured magnetic parameters are schematically illustrated

in a hysteresis loop (Fig. 1.4). The soft magnetic materials possess narrow hysteresis loop

with low coercive field strength [18]. The magnetization also increases with the increase

in magnetic field strength. In an external magnetic field large enough, the spin in each

domain rotates parallel to the direction of applied magnetic field till the proper alignment

of all dipoles. Thereafter, the magnetization flattened at a value called the saturation

magnetization (MS) (Fig. 1.4). As the strength of the magnetic field decreases, spins cease

to be aligned with the field leading to the decrease in total magnetization.

In ferrimagnets, some magnetic moment remains at zero fields in the form of

residue. This value of residual magnetization is called the remanent magnetization (Mr).

The coercive field, HC is the magnitude of the field that should be applied in the negative

direction to bring the magnetization of the sample back to zero. The shape of hysteresis

13

loop is of particular interest in magnetic recording applications that should be (ideally) a

square hysteresis loop with a large magnetization and moderate coercivity [28].

Figure 1.4: The hysteresis loop of a magnetic material, where H is the magnetic field

amplitude and M is the magnetization of the material [18].

1.4.1.2 Magnetic anisotropy

In most of magnetic materials, the magnetization relatively tends to align itself

along an easy axis of magnetization. All ferromagnetic and ferrimagnetic materials

exhibit, to a lesser or greater degree, a crystal direction or a set of directions along which

the magnetization prefers to be occurring [10]. The most important property of magnetic

materials is magnetocrystalline anisotropy. The magnetocrystalline anisotropy of the

materials has much association with their crystal symmetry.

There are three situations that lead to this anisotropy as an intrinsic property of

crystalline materials. The most important is that in which the atoms have an electron-

orbital moment along with an electron-spin moment. In this situation, the spin direction

can be coupled to the crystal axis. This occurs through both coupling between the spin

and orbital moments and the interaction between the charge distribution on the orbit and

14

the electrostatic field created by surrounding atoms. There will then be one or more axes

or surfaces along which magnetization is more feasible with little work. The crystal will

then be preferably magnetized along such an easy axis or plane. The second situation

occurs in non-cubic crystal lattices. In this situation, the magneto-static interaction

between the atomic moments is also anisotropic, which can lead to easy magnetization

directions or planes. The third possibility of crystalline anisotropy is in the directional

control of atoms as described by Néel elsewhere [29].

1.4.2 Electrical properties

1.4.2.1 DC-electrical properties

The electrical properties in ferrites depend upon a number of factors such as;

synthesis route, annealing temperature, amount and type of substituents, etc. Ferrites have

much higher electrical resistance than metals and they are regarded as very structure-

sensitive materials, their resistivity at room temperature can vary between 10-2

-1011

Ωcm

[30]. The electrical resistivity can further be increased by special sintering conditions and

selecting suitable compositions with small amount of metal oxides substitution in ferrites

[31].

The conductivity mechanism in crystalline materials depends upon by the transfer

of either electron or ion. Generally conduction by one or more type of charge carrier is

prominent, but in some inorganic materials both ionic and electronic conductions

predominate [32]. According to the Verwey model [33] of conduction, the electronic

conduction in ferrites is mainly due to the hopping of electron between ions of the same

element having different valence state at octahedral sites. The hopping probability

depends directly upon the distance between ions involved in conduction and the potential

barrier that have to be overcome known as activation energy [34]. Mg-ferrite is a hopping

or a small polaran hopping semiconductor with a large concentration of mobile electrons

[35]. The conduction process in Mg-ferrite is mainly due to electronic hopping between

the Fe3+

and Fe2+

ions, placed at the adjacent octahedral sites [34]. This exchange process

depends upon the mutual overlapping of orbital of transition metal ions and oxygen [36].

15

The activation energy of hopping can be calculated from the temperature-

dependence of resistivity (lnρ vs. 103/T) and values are found to be in the range of 0.3-

0.7 eV for spinel ferrites, as also reported by Iqbal et al [34, 37] for Co-Cr and Ni-Cr

doped Mg-ferrites. The high values activation energy (0.4-0.6 eV) approves the high

resistivity of the compound at room temperature. In addition, these values are larger than

the transition energy of Fe2+/Fe3+ ion-pair (Ee = 0.2 eV), which indicates that the small-

polaron model of electron-hopping is favored in the studied samples [38]. Small polaron

is defined as a combination of hopping electron and its accompanying lattice polarization

field [39]. The hopping process occurs between eg orbitals on adjacent Fe3+

/Fe2+

cations

on the octahedral lattice site [40]. So the hopping and hence the resistivity in ferrites

depends on the number of ferrous and ferric ions in the material.

1.4.2.2 Dielectric properties

Ferrites have taken much attention as dielectric materials, owing to their high

electrical resistivity and low dielectric losses. An important feature of a dielectric is its

ability to support an electrostatic field to reduce the dissipation of energy in the form of

heat. The smaller the dielectric loss, the more useful is a dielectric material. Another

element is the dielectric constant, the degree to which a substance concentrates on the

electrostatic lines of flux. In the classical approach to the dielectric model, a material is

made up of atoms. In the presence of an electric field, the cloud of negatively charged

electrons of an atom in dielectrics is distorted, creating dipoles as shown in the left of the

Fig. 1.5. As each dipole is characterized by its dipole moment thus they produce their

own field, which interact with the external applied field [41]. The process of relative

displacement of the negative and positive charges of atoms or molecules, the orientation

of existing dipoles toward the direction of the field, or the separation of mobile charge

carriers at the interfaces of grain-boundaries, caused by an external electric field is

referred to as an electric polarization [42].

There are four types of polarization, as electronic, Ionic or atomic, dipolar and

interfacial or space-charge polarization. Electric polarization associated with mobile and

trapped charges is generally referred to as interfacial or space-charge polarization. This

mainly occurs in amorphous or polycrystallin

carriers (electrons, holes or ions)

Figure 1.5: Electric field interactions with an atom under the classical dielectric model

[41]

One of the important electrical properties of dielectric materials is permittivity or

dielectric constant. The dielectric constant depends strongly on the frequency of the

alternating electric field or the rate of the change of the time

depends upon the chemical structure and the imperfections (defects) of the material, as

well as on other physical parameters including temperature and pressure, etc.

Spinel Mg-ferrite and its allied compounds

microwave and surface mount devices due to their high electrical resistivity and low

values of dielectric constant and dielectric losses.

1.5 Advantages o

Most of technologically

exhibit low electrical resistivity,

transformer cores, inductor cores and other applications

Mg-ferrite and its metal doped derivatives

materials because of their

mainly occurs in amorphous or polycrystalline materials consisting of traps and

carriers (electrons, holes or ions) [43].

Electric field interactions with an atom under the classical dielectric model


. The dielectric constant depends strongly on the frequency of the

alternating electric field or the rate of the change of the time-varying

on the chemical structure and the imperfections (defects) of the material, as

well as on other physical parameters including temperature and pressure, etc.

ferrite and its allied compounds might be useful


values of dielectric constant and dielectric losses.

Advantages of Magnesium Ferrite

technologically important magnetic materials like iron and metallic alloys

exhibit low electrical resistivity, making these materials useless for applications i

transformer cores, inductor cores and other applications operating

ferrite and its metal doped derivatives have advantage over other electro

their high inherent resistivity, high permeability and stability over a

16

e materials consisting of traps and charge

Electric field interactions with an atom under the classical dielectric model


. The dielectric constant depends strongly on the frequency of the

varying fields. It also

on the chemical structure and the imperfections (defects) of the material, as

well as on other physical parameters including temperature and pressure, etc., [41].

be useful for applications in


iron and metallic alloys

useless for applications in

at higher frequencies.

over other electro-magnetic

high permeability and stability over a

17

wide range of temperature. With these advantages, substituted Mg-ferrites outweigh all

other magnetic materials [44].

The trouble with the other electro-magnetic materials is their low electrical

resistivity which offers the induced currents (called eddy currents) to flow within the

materials, thus producing heat. This generated heat is often a serious problem and

become a cause for wastage of energy. Thus, the materials become inefficient due to

wastage of energy, especially as the frequency is high enough. However, the performance

of soft magnetic materials like Mg-ferrite is much better at high frequencies due to their

high electrical resistivity [45-46].

Moreover, high permeability and time/temperature stability are some additional

important features that have expanded the usage of soft ferrites in high frequency and

delay lines, broadband transformers, adjustable inductors, quality filter circuits and other

high-frequency circuit’s electronics. At high frequencies, the use of soft ferrites is

relatively more systematic with respect to that of the other circuit components whose

performance must be improved. An important factor in the choice of soft ferrites is that

they are generally less expensive than magnetic metals and alloys. Soft ferrites are the

best option of the core material for operating frequencies from 10 kHz to a few hundred

MHz with proper combination of low cost, high inductor quality, high stability and low

volume. In addition, no other magnetic material possesses magnetic and mechanical

parameters as flexible as those of soft ferrites.

1.6 Applications of Magnesium Ferrite

Crystalline magnesium ferrite is a soft magnetic material [47] and it is an

important member of the spinel family. It is n-type semiconducting material having

number of applications in adsorption, sensors, and in magnetic technologies [47]. In

nanocrystalline magnesium ferrite, many of the useful properties of its crystalline

counterpart, such as magnetization, are enhanced. Further, magnesium ferrite belongs to

class of soft magnetic materials which is easy to magnetize and demagnetize, so are used

in electromagnets. Owing to its nanocrystalline nature and useful properties, the material

18

shows a good potential for novel applications in humidity, gas sensing [48] and drug

delivery [11, 49]. MgFe2O4 is also known for its good photoelectric effect [47, 51-52].

Apart from its magnetic and electronic applications, MgFe2O4 finds a number of

applications in heterogeneous catalysis [53-55]. Moreover, magnesium ferrite and its

allied compounds have widespread applications in microwave devices such as circulators,

insulator, phase shifters and multifunctional devices [56-58] because of their low

magnetic and dielectric losses and high resistivity.

1.7 Synthesis and Characterization of Magnesium Ferrite: Literature

Review

Oliver et al. [59] used the sol-gel supercritical drying method to synthesize the

fine powder of MgFe2O4, which was calcined at two different temperatures i.e. 773 K and

1073 K. The powder structure was matched with the spinel phase of MgFe2O4 and α-

Fe2O3, as an impurity been observed in samples. Superparamagnetic nature was observed

in as-produced powders at room temperature, having single magnetic domain particle

with size of 11 nm. The particle size distribution was evaluated by fitting the

magnetization data to a Langevin function, and confirmed by Mossbauer spectra

measured at temperature range 25-298 K. Very similar particle size distributions were

observed for all three methods. The saturation magnetization and magnetocrystalline

anisotropy values were observed to be same for both bulk and nano-particle and did not

change with the size of particles.

Bhosale et al. [60] synthesized mixed Cu-Mg-Zn ferrites by coprecipitation

method. The lattice parameter decreases gradually while density increases with the

increase in Mg2+

substitution level. This variation is attributed to smaller ionic size of

Mg2+ ions. The initial permeability (µ i) is found to be affected with magnetization (MS)

and particle size of (D) samples. The µ i decreased with increasing Mg2+

content due to a

lower value of the anisotropy constant (Ki) for MgFe2O4 than that for CuFe2O4.

Zimnol et al. [61] prepared the epitaxial thin films of MgFe2O4 by solid state

reactions between MgO substrates and FeO vapors. The different compositions of

19

epitaxial spine1 films were obtained having wide range of magnetic properties. The

deposited films were characterized by using X-ray diffraction (XRD), Rutherford

backscattering spectroscopy (RBS), energy dispersive X-ray spectroscopy (EDX),

transmission electron microscopy (TEM) and selected area electron diffraction (SAED).

Magnetic hysteresis loops were measured using the magneto-optical Faraday Effect.

Kuznetsov et al. [62] synthesized MgFe2O4 and chromium substituted

magnesium-zinc ferrite Mg0.5Zn0.5Fe2-xCrxO4 (0≤x≤1.5) by self-propagating high

temperature synthesis (SHS), a combustion process involving the reaction between

magnesium, zinc, iron and chromium oxides with iron or chromium metal powders and

sodium perchlorate. Two series of SHS samples were produced with or without magnetic

field of 1.1 T followed by annealing at 1400 °C for 2 h. The pure cubic phase of spinel

ferrite was produced in both cases as observed by X-Ray data. Zn and Cr contents affect

the cubic lattice parameter. Mössbauer studies at 298 and 80 K showed a remarkable

change in sublattice occupancy with Cr content. Magnetic data showed that the coercivity

of doped samples is higher than that of pure composition. The use of a magnetic field was

found to influence the microstructure and magnetic properties of material.

Chandrasekaran et al. [63] prepared the mixed Mg1-xZnxFe2O4 ferrites with

different sintering conditions as compare to previous reports. The samples were

ferrimagnetic with multidomain grains as indicated by AC magnetic susceptibility and

hysteresis studies.

Chen et al. [64] prepared the MgFe2O4 nanoparticles by coprecipitation method.

Magnetic measurements and neutron diffraction have shown the existence of a

superparamagnetic state in the synthesized system. The superparamagnetic relaxation was

studied by using Mössbauer spectroscopy and relaxation time has been correlated with

the particle size and temperature which is consistent with Neel theory.

Ravinder et al. [65] studied the dielectric properties of mixed Mg-Zn ferrites in

the frequency range of 1-100 kHz using a capacitance bridge. The dielectric constant vs.

frequency curve shows a normal dielectric behavior of ferrite materials. The frequency

dependence of a dielectric loss tangent (tanδ) curve possessed a peak at a certain

20

frequency for all the prepared compositions. The dielectric constant for these mixed

ferrites is almost inversely proportional to the square root of DC-resistivity. These

observations can be explained on the basis of an electron hopping between Fe2+ and Fe3+

ions.

Chen et al. [49] prepared MFe2O4 (M=Cu, Zn, Cd and Mg), with a large specific

surface area of 40-80 m2/g by a coprecipitation method and tested for sensing the

reducing gases like CO, H2, LPG, C2H5OH and C2H2. The gas sensitivity of ferrites

varied in the order MgFe2O4~CdFe2O4>CuFe2O4>ZnFe2O4 due to varying the specific

surface area or grain size. It was observed that MgFe2O4 and CdFe2O4 were the most

sensitive and selective to LPG and C2H2, respectively, among the ferrites tested.

Liu et al. [66] studied a correlation between the electron spin-orbital angular

momentum coupling and the superparamagnetism in MgFe2O4 and CoFe2O4

nanoparticles. The neutron diffraction studies have shown that the cation distribution and

contribution to the magnetic anisotropy from the Fe3+ lattice sites is almost the same in

both nanocrystallites. The blocking temperature of CoFe2O4 nanoparticles is 150 oC

higher than that of the same sized MgFe2O4 nanoparticles due to more anisotropic nature

of Co2+

ions and confirmed by Mössbauer spectroscopic studies which demonstrate that

the magnetic anisotropy of CoFe2O4 nanoparticles is higher than that of the same size

MgFe2O4 nanoparticles, which can be controlled by adjusting the magnetic anisotropy

energy of nanoparticles.

Sěpelak et al. [67] investigated the effects of high energy milling on MgFe2O4.

The crystallite size of MgFe2O4 can be reduced to nanometer range by mechanical

treatment and also control the redistribution of cations between tetrahedral and octahedral

sites. The thermal stability range of mechanically induced metastable states is studied by

the change in temperature.

Men et al. [68] re-examined the experimental results obtained during reduction of

oxide solutions because of no proper interpretation previously. Some approximations

were made in order to describe the reduction of systems as complicated as MIx M

II1-x

MIII

yFe2-yO4 (MI = Mg

2+, Mn

2+; M

II = Fe

2+, M

III = Cr

3+; Al

3+; V

3+). New schemes for

21

calculating phase equilibrium provided more adequate interpretation of the early

reduction data.

Rana et al. [69] prepared a series of spinel Mg1-xNixFe2O4 (x = 0.0-1.0) using

standard ceramic method. AC susceptibility measurements were performed to calculate

lande-g factor (g), effective magnetic moments (Peff), Curie temperature (TC),

paramagnetic Curie temperature (ө (K)) and exchange integral (J/k). The g-values, Peff

and θ(K) show increasing trend with the increase in Ni content up to x = 0.75 while TC

continues to decrease. The decreasing trend in magnetic moments (nB) and θ(K) for

x>0.75 could be correlated to the distribution of cations on A and B sites. The Y-K angle

for NiFe2O4 shows a gradual increase with Ni contents and approaches to 600.

Qi et al. [70] prepared Mn substituted MgCuZn ferrites (Mg0.2Cu0.2Zn0.6O) (Fe2-

xMnxO3)0.97 (x = 0.00-0.07) using nanosized precursor powders synthesized by a sol-gel

method. It has been observed that MgCuZn ferrites doped with Mn possess higher initial

permeability and better grain structure than that of NiCuZn ferrites prepared by the same

method could be ideal materials for high inductance multilayer chip inductor. The

variation of initial permeability of MgCuZn ferrites with the Mn substitution might be

attributed to the decrease of magnetostriction constant.

Antoshina et al. [71] investigated the magnetic properties of CuGaxAlxFe2-2xO4 (x

= 0.2-0.7), CuGaxAl2xFe2-3xO4 (x = 0.1-0.5), CuxNi1-xFe0.6Cr1.4O4 (x = 0.0-0.4), CuFe2-

xCrxO4 (x = 0.0-1.6) and of diluted ferrites-chromates GaxFe1-xNiCrO4 (x = 0.0-0.8). It is

revealed that the effective magnetic anisotropy constant of frustrated magnetic structure

is lesser in magnitude than that for ordered ferrites. The transition from anisotropic

magnetic state to isotropic state takes place at temperature Tt, smaller than Curie

temperature TC in frustrated structure ferrites.

Sěpelak et al. [72] investigated the changes in MgFe2O4 caused by high-energy

milling, using Mössbauer spectroscopy, magnetization measurements, and electron

microscopy. The observed enhancement of the magnetization in nanoscale milled

MgFe2O4 is discussed on the base of cation redistribution and spin canting.

22

Radwan et al. [73] prepared the Mg1+xTixFe2-2xO4 of single phase spinel structure

assured by X-ray analysis. The octahedral preference of Mg2+

ions doped on the expense

of the Fe3+ ions increases the inversion factor of spinel. The conduction is occured due to

mobility of thermally activated charge carriers. The mobility of charge carriers could be

explained using Verwey model of conductivity which based on the electron hopping

between Fe2+

/Fe3+

located on the same sub-lattice sites. The octahedral site preference of

Ti4+

ions decreases the conductivity of the samples. Unusual behavior of Ti content of 0.7

and 0.8 was observed due to the presence of impure phases.

Rabanal et al. [74] studied the magnetic properties of powdered Mg-ferrite

processed with a centrifugal mill. The starting ferrite powder was prepared by solid-state

reaction occurred at 1400 oC for 48 h. The crystalline size and internal strain were

evaluated by XRD data using Williamson-Hall and Debye-Scherrer methods. The

nanoparticles were obtained for low milling time. The X-ray analysis indicates that the 17

h of milling caused the appearance of two α-Fe2O3 peak. The saturation magnetization

remains nearly constant at 39.2 emu/g indicates the lack of inversion degree even for

longer milling times, while coercivity increase up 576.7 Oe due to internal stresses

caused by the mechanical grinding.

Chauhan et al. [75] used citrate precursor method to modify the magnetic

properties of substituted Mg-Mn ferrites successfully. The citrate precursor method is

used to obtain close composition control, better homogeneity, lower processing

temperature, higher purity, lower porosity and more uniform grain growth. The various

magnetic parameters like saturation magnetization, magneton number, and thermal

variation of AC-susceptibility, Neel temperature and initial permeability were calculated.

The saturation magnetization and initial permeability of the synthesized materials were

improved significantly with indium and cobalt substitution. The Neel temperature of the

samples was increased with the increase in cobalt content. Magnetic losses were 1-2

orders lower as compared to those for Mg-Mn ferrites prepared by conventional ceramic

method, suitable for high-frequency applications.

Chhaya et al. [76] investigated the Ca2+

doped MgFe2O4 without altering the

cubic symmetry and affecting the structural, magnetic and electrical properties of Mg1-

23

xCaxFe2O4 (x = 0.0-0.35) spinel ferrite system studied by means of X-ray diffraction

(XRD), magnetization, AC susceptibility and DC resistivity measurements. The variation

of magneton number with Ca2+ content can be explained on the basis of Néel’s collinear

spin model. The Néel temperature determined through AC susceptibility and DC

resistivity measurements is in close agreement with theoretical values. The variation in

electrical resistivity coincides with the change in activation energy.

Turkin et al. [77] synthesized MgFe2O4 annealed at 1373 K using silica-tube

technique. The product is homogeneous and fine grained with the inversion parameter of

0.75. The calorimetric measurements indicate the second-order phase transition of

antiferromagnetic to paramagnetic materials at 597 K. This silica-tube technique of the

synthesis prevents the escape of oxygen at long heating.

Verma et al. [78] synthesized nanosized magnesium ferrite using mild microwave

hydrothermal (MH) conditions. The average particle size of the ferrite is found to be ~3

nm calculated from X-ray diffraction and transmission electron microscopic analyses.

Vibrating sample magnetometric studies indicate the superparamagnetic nature of ferrite

particles.

Thummer et al. [79] conducted 57Fe Mössbauer spectroscopy at 300 K to

investigate the magnetic behavior of the spinel ferrite, MgAlxCrxFe2-2xO4 (x = 0.0-0.5).

The Mössbauer spectra for x = 0.0-0.2 exhibit two Zeeman sextets due to Fe3+

ions

distributed over two sites i.e. octahedral and tetrahedral sublattice sites, while a central

paramagnetic doublet superimposed on the magnetic sextets appeared for x = 0.3-0.5. He

discussed the variation of hyperfine interactions and the appearance of the central doublet

on the Zeeman sextet thoroughly.

Liu et al. [80] synthesized n-type nanomaterials (MgFe2O4) by convenient,

environment friendly, inexpensive solid-state reaction method. The material structure and

crystallite microstructure of samples have been evaluated by X-ray diffraction (XRD),

transmission electron microscopy (TEM) and high-resolution transmission electron

microscopy (HRTEM). Conductance responses of the MgFe2O4 were measured by

exposing the thick film to reducing gases like methane (CH4), hydrogen sulfide (H2S),

24

liquefied petroleum gas (LPG) and ethanol gas (C2H5OH) and observed various sensing

responses to these gases at different operating temperature. Successive on and off

responses have been repeated and no major changes in the response signal were seen.

Pradeep et al. [81] prepared nano-particles of M0.5Mg0.5Fe2O4 (M=Ni, Cu and Zn)

using the sol-gel method. Synthesis of single phase polycrystalline ferrite materials is

assured using XRD. Lattice constant and particle size have been determined from XRD

data. Mixed CuZn-ferrites have greater values of the lattice constant as they are bigger

ions than Ni. Particle size was decreased by substitution of Cu to Zn. The samples were

subjected to VSM measurements and FT-IR characterization. The magnetic moment

values are determined and the theoretical calculation was done to propose the cation

distribution. FTIR data were analyzed for the respective sites. The higher frequency band

(v1) and lower frequency band (v2) are attributed to the tetrahedral and octahedral

complexes, respectively. The difference in the trends of bond length and force constant is

able to elucidate the role of crystal field effect.

Lakshman et al. [82] prepared the Mg0.9Mn0.1InxFe2-xO4 and Mg0.9Mn0.1CryFe2-yO4

by the conventional ceramic route. The influence of In3+

and Cr3+

ions on the DC

resistivity, dielectric constant and dielectric loss factor are evaluated in this paper. The

resistivity increases with the increase in dopant contents. The activation energy and

dielectric constant were found to increase with the substitution level of In3+ and Cr3+ ions.

The dielectric loss tangent (tan δ) measured at 100 kHz and 13 MHz are found to be very

small for the samples with higher dopant contents. The smaller values of loss factor

indicate that the prepared materials might have good potential for microwave devices.

Huang et al. [52] synthesized the nanocrystallites of MgFe2O4 by a sol-gel

combustion method. A well-crystalline MgFe2O4 was obtained by annealing at 500 oC for

2 h. The synthesized MgFe2O4 were investigated by TG-DSC, FT-IR, XRD and TEM

analyses. The average particle size was found to be 10 nm with a narrow size distribution.

The particle size increased with the increase of annealing temperature. The MS of the

MgFe2O4 annealed at 500-900 oC are increased from 9.9 to 30.6 emug

−1, respectively.

25

Masti et al. [83] studied the magnetization and permeability of polycrystalline

materials, CdxMg1-xFe2-yCryO4 (x = 0.0-1.0; y = 0, 0.05 and 0.10). The Neel’s two-

sublattice model exists up to x = 0.4, for y = 0, 0.05 and 0.1 and a three-sublattice model

(YK-model) is predominant for x>0.4 and y = 0, 0.05 and 0.10 revealed by the

magnetization. The MS was found to decrease with the increase in Cr3+

contents, which is

attributed to the weakness of B-B site interaction. Variation of initial permeability with

temperature revealed the long-range ferromagnetic ordering in the compounds with x =

0.4. The sample with the composition x≤0. 4 and y = 0, 0.05 and 0.10 possess peaking

behavior near the Curie temperature attributed to the decrease of anisotropy constant K1

to zero. Addition of Cd2+

and Cr3+

resulted in a decrease in Curie temperature and initial

permeability, respectively.

Modi et al. [84] studied the elastic behavior of MgAlxFe2-xO4 (x = 0.0-1.0) by IR

spectroscopy and X-ray diffractometer. The IR spectra and X-ray analyses are used to

determine the force constants for A- and B-sites and calculated lattice constant, elastic

moduli-like bulk modulus, Young’s modulus, rigidity modulus, Poisson’s ratio and

Debye temperature. The observed variation of elastic constants has been correlated to the

strength of interatomic bonding and electronic configuration of the cations involved in

the system. The applicability of the heterogeneous metal mixture rule has been tested to

find the validity of results obtained from the present method and compared with the

results of the other method.

Bergmann et al. [85] reported the single-step synthesis of nanosized MgFe2O4 via

mechano-chemical processing of oxide precursors. The synthesized materials were

subjected for X-ray diffraction and 57

Fe Mössbauer spectroscopic analyses. The

transmission electron microscopy assured the nanoscale nature of the mechano-

synthesized material.

Lakshman et al. [86] studied the effect of the substitution of Cr3+

in mixed

MgCuMn ferrites, Mg0.9Cu0.1Mn0.05CrxFe1.95-xO4 (x = 0.0-0.9) with incorporation of small

amounts of Cu2+

ions. X-ray diffraction and Mössbauer spectroscopic analyses of these

samples have been performed. Mössbauer analysis at 300 K showed two Zeeman sextets

for lower concentration of Cr3+

followed by relaxation phenomenon and the spectrum

26

shows paramagnetic doublet for x =0.9. The variation in Mössbauer parameters, viz,

isomer shift, quadrupole splitting and hyperfine magnetic field with dopant concentration

has been evaluated.

Kong et al. [87] studied the densification, grain growth, and microstructure of

Mg-Cu-Co ferrite (MgFe1.98O4, Mg1-xCuxFe1.98O4) with x = 0.10-0.30 and Mg0.90-

xCoxCu0.10Fe1.98O4) with x = 0.05-0.20 fabricated for high frequency antennas.

Magnesium ferrite (MgFe1.98O4) is a promising magneto-dielectric material for high

frequency antennas. But due to its poor densification, it could not be sintered at a

temperature below 1200 oC. High-temperature sintering resulted in a drastic but

undesirable change in electrical and dielectric properties. The poor densification and slow

grain growth rate have been improved by the addition of Cu ions into MgFe1.98O4. While,

the presence of Co did not have any significant influence on the densification and grain

growth rate of the synthesized system.

Sěpelak et al. [88] reported the magnetization enhancement in nanosized

mechano-synthesized MgFe2O4 and discussed on the basis of a modified core-shell

model. Due to random distribution of cations, the surface shell of nanoparticles exhibits

an effective magnetic moment 2 times greater than that of the particle core. The thickness

of surface shell is calculated from the size-dependent magnetization measurements.

Ichiyanagi et al. [89] prepared the MgFe2O4 nanoparticles by using wet chemical

method. The particle size estimated from X-ray diffraction patterns were in the range of 3

to 8 nm. Magnetization measurements were performed under an applied field of 750 kOe.

Both the field-cooled (FC) and the zero-field-cooled (ZFC) magnetization dependent

blocking temperature, Tb, were observed to be around 30 K. A high coercivity of ~1000

Oe was observed at 5 K. A big difference in magnetization was observed between the

quenched samples and annealed samples.

Deng et al. [90] developed a low-temperature solvothermal synthetic route for

synthesis of uniform-sized MFe2O4 (M = Mg, Cu, Ni) ferrite. The size of the as-prepared

ferrite microspheres could be controlled in the range of 300–800 nm by adjusting some

growth parameters. The X-ray powder diffraction (XRD), transmission electron

27

microscopy (TEM), scanning electron microscopy (SEM) and energy-dispersive X-ray

analysis (EDAX) were employed for structural investigation of the as-prepared ferrite

microspheres. The magnetic properties measurements were also carried out for as-

prepared materials.

Vital et al. [91] prepared the nanoscale ZnFe2O4, Mg0.5Zn0.5Fe2O4 and

Mg0.2Cu0.2Zn0.62Fe1.98O3.99 powders by flame spray synthesis (FSS). Particle size was

estimated in the range of 6–13 nm. Compacts prepared from Mg–Cu–Zn ferrite

nanoparticles possessed an extremely high sinter-activity. A sintered density of 5.05 g

cm-3

was achieved after sintering for 2 h at 900 oC which is greater than that achieved

(4.91 gcm-3

) from the conventional ceramic route. The permeability of the sintered Mg-

Cu-Zn ferrite compacts was reached to value of µ = 600 at 1 MHz and MS was 80 emug-1.

The significant sintering activity of the synthesized ferrite powders is attributed to their

small particle size.

Shah et al. [92] reported the addition of 2 wt%, 4 wt% and 6 wt% cerium oxide to

MgFe2O4 prepared by a ceramic method. They observed the decrease in resistance from

52MΩ to 2MΩ with the addition of 6 wt% cerium oxide. The humidity generated is in the

range of 10% RH to 90% RH at 25 C. The addition of cerium oxide increases the

intergranular porosity, distribution of pore size and open pores in the system. The

addition of cerium oxide improves the sensitivity and shows a better linearity than that of

pure magnesium ferrite. X-ray diffraction (XRD) analysis confirms the spinel phase

formation of the samples and the calculated value of porosity is confirmed by the

scanning electron micrographs of the samples.

Sharma et al. [93] studied the magnetic properties of nanosized

Mg0.95Mn0.05Fe2O4 samples and also characterized by X-ray diffraction, Mössbauer

spectroscopy, DC magnetization and frequency dependent real χ′(T ) and imaginary χ″(T

) parts of AC susceptibility measurements. Mössbauer measurements show a magnetic

transition to an ordered state at 195 K. The ZFC curve shows a broad maximum at Tmean

= 195 ± 5 K, which depends upon the distribution of particle volumes in the sample. A

frequency-dependent peak is well described by Vogel-Fulcher law, giving a relaxation

time τo = 5.8 × 10-12

s and an interaction parameter To = 195 ± 3 K. These values are

28

evident in strong interactions between the nanoparticle system. On the other hand fitting

with the Néel-Brown model and the power law provide an unrealistic large value of τo

(~6×10-69 and 1.2×10-22 s respectively).

Pradeep et al. [94] synthesized MgFe2O4 nanomaterials using sol-gel

autocombustion method and performed the X-ray diffraction analysis for structural

studies of the as-synthesized materials. The particle size and lattice constant are obtained

using XRD data. The site preference of cations in nano-MgFe2O4 is compared with its

bulk counterpart. The Scanning Electron Microscopy (SEM) was utilized for

morphological features of the prepared nanoparticles. Fourier transform infrared

spectroscopy (FTIR) assures the formation of spinel structure and revealed the effect of

nanosized synthesis on various parameters such as bond length, vibration frequency and

force constant. The M-H loops of specimens have been measured using vibrating sample

magnetometer (VSM) and magnetic parameters such as saturation magnetization (MS),

coercivity (HC) and retentivity (MR) are determined from VSM data.

Li et al. [95] fabricated a jingle-bell-shaped hollow structured Ag@MFe2O4 (M =

Ni, Co, Mg, Zn) nanomaterials by direct adsorption of metal cations Fe3+

and M2+

on the

surface of the Ag@C spheres followed by calcination to remove the middle carbon shell

and transform the metal ions into pure phase ferrites. The as-prepared composites were

subjected for characterization using X-ray photoelectron spectroscopy, energy-dispersive

X-ray analysis, X-ray powder diffraction, scanning electron microscopy, transmission

electron microscopy, UV-vis spectroscopy and SQUID magnetometer. The observed

results showed that the composites took the magnetic character of the ferrite shell and

optical together with antibacterial characteristics of the Ag core.

Hankare et al. [24] reported that the mixed metal oxides with spinel structure

possessed interesting structural, magnetic, electrical, and catalytic properties. Chromium

substituted magnesium ferrites were prepared by co-precipitation technique and subjected

to investigation. The spinel phase formation was assured by XRD analysis. All the

samples showed decreasing trend of lattice constant from 8.40 to 8.33Å with an increase

in dopant content. Two strong bands are evident from infrared spectra, one around 600

cm-1

is attributed to the intrinsic vibrations of tetrahedral complexes and second one

29

appeared at 400 cm-1

is because of octahedral complexes. The morphological features

were investigated by scanning electron microscope while element analysis was carried

out by dispersive X-ray spectroscopy. The synthesized materials (MgFe2-xCrxO4) were

also investigated for thermal, electronic and magnetic studies.

Barati [96] synthesized the mixed ferrite with compositions Mg0.80-

xCu0.20ZnxFe2O4 (x = 0.5, 0.55, 0.60 and 0.63) through nitrate-citrate gels auto-

combustion method. TG/DTA analysis was carried out to study the combustion process

of the dried gels. The obtained precursor was calcined at 800 oC for 1 h followed by

sintering at 900 oC for 4 h. The obtained materials were characterized for phase

identification, grain size and lattice parameter determination using X-ray diffraction. The

magnetic and electrical properties of synthesized materials have also been investigated

for various parameters. The initial permeability, saturation magnetization, dielectric

constant and dielectric loss were found to increase and AC-resistivity was decreased with

the increase in dopant contents. The prepared material is suitable for the application in

multilayer chip inductor due to its good magnetic properties and low loss at high

frequency.

Gadkari et al. [97] investigated the structural properties of Sm3+

doped Mg1-

xCdxFe2O4 (x = 0.0-1.0) synthesized by oxalate co-precipitation method. The samples

were annealed at 1050 °C for 5 h. The cubic spinel structure was assured by X-ray

diffraction analysis and evaluated the different parameters like lattice constant, X-ray

density, physical density, porosity, crystallite size, site radii and bond length on

tetrahedral and octahedral sites. The lattice constant increases with an increase in Cd2+

content and crystallite size varies from 28.69 to 32.05 nm. The surface morphology

shows an increased grain size of the samples with an increase in Cd2+ contents. The IR

spectra show two strong absorption bands at 5.87×104 m

-1 and 4.27×10

4 m

-1 for

tetrahedral and octahedral sites, respectively.

Gaudon et al. [98] studied the relationship of cationic and anionic vacancies

and/or interstitial atoms instead of a partial inversion rate of the A2+

and B3+

cations for

(Zn2+

/Mg2+

)[Fe3+

]2O4 spinel oxides. They observed the veracity of the new structural

formalism by using the X-ray diffraction patterns and Mössbauer spectroscopy.

30

Hankare et al. [99] reported the mixed metal oxides exhibiting the interesting

structural and electrical properties. The MgFe2-xCrxO4 was synthesized by the simple

coprecipitation method and has been investigated by X-ray diffraction and scanning

electron microscopy. The IR spectrum shows two strong bands at 600 and 500 cm-1

. Both

n-type and p-type behaviors were observed by the thermo-electric power measurements

carried out from room temperature to 500 °C.

Chand and Singh [100] have synthesized the MgGd0.1Fe1.9O4 by conventional

ceramic method with enhanced electrical and dielectric properties. X-ray analysis assured

the formation of single-phase structure. The DC-resistivity is enhanced by one order of

magnitude with Gd substitution in Magnesium ferrite. High resistivity along with lower

dielectric loss is attributed to better compositional stoichiometry and nature of the

additives. Dielectric properties of the sample have been studied in the frequency range of

0.1-20 MHz at various temperatures. Possible mechanisms correlated with the results

have been discussed shortly in this paper.

Sasaki et al. [101] synthesized magnesium ferrite (MgFe2O4) by hydrothermal

route for which suspensions of Mg(OH)2 and Fe(OH)3 in appropriate ratio was prepared

and pressurized to 30 MPa by high-pressure pump followed by rapid heating to the

reaction temperature. The Mg/Fe molar ratio varied from 0.5 to 1.5 to obtain single-phase

MgFe2O4. The stoichiometric ratio of Mg/Fe = 0.5 contains both MgFe2O4 and α-Fe2O3

while, at Mg/Fe = 1.0 and 1.5, the single-phase MgFe2O4 is obtained. MgFe2O4

synthesized in the present study with particle size of 20 nm, exhibits superparamagnetic

behavior.

Ahmed et al. [102] synthesis magnesium ferrite compacts by using a mixture of

magnesite ore and waste iron oxide (mill scale). They investigated the effect of different

mixture composition and sintering conditions on the phase change, and then evaluated the

compressive strength, physical and magnetic properties of sintered compacts. A single

phase ferrite obtained from a mixture of 40 weight% magnesium ore and 60 weight%

mill scale possessed low porosity and high saturation magnetization.

Farghali et al. [103] studied the variation of physical characteristics by embedding

the ferrite materials into polymeric matrices. They prepared well dispersed

polyaniline/Co1-xMgxFe2O4 nano-composite (x = 0, 0.5, 1) with good magnetic and

31

electrical properties. TGA results indicated that the ferrite nano-particles could improve

the thermal stability of composite. The electrical conductivity of the pure polyaniline

decreased while the saturation magnetization (MS) and coercivity (HC) increased by

embedding the ferrite nanoparticles in composite.

Burianova et al. [104] presented a comparative study of magnetocaloric effect

(MCE) in superparamagnetic (SPM) regime in two different types of magnesium viz;

MgFe2O4 encapsulated in amorphous SiO2, or as matrix-less nanoparticles synthesized in

supercritical water conditions. XRD and TEM analyses were employed to assess the

particle diameter of all prepared samples. The entropy change, ∆S was derived from the

magnetization, M(H,T) curves measured at specific temperature intervals. They observed

a broad peak of ∆S for all the samples in the temperature range above the TB.

Reddy et al. [105] prepared the polycrystalline MgCuZn-ferrites with chemical

composition Mg0.50-xCuxZn0.50Fe2O4 (x = 0. 00-0.30) by the microwave sintering process.

The precursors were calcined and sintered at 950 °C for 30 min. Structural, elemental and

morphological analyses were performed using X-ray diffraction, energy dispersive X-ray

spectrometry and scanning electron microscope, respectively. The lattice parameter was

observed to increase with the increase in copper content. A significant densification is

observed with the incorporation of Cu ions in the synthesized materials. The prepared

ferrites were analyzed for initial permeability, dielectric constant and loss tangent and AC

conductivity measurements. The dielectric properties were measured in the frequency

range of 100 Hz to 1 MHz. Initial permeability and dielectric constant were observed to

increase, while dielectric loss tangent decreased with Cu substitution up to x = 0.20. The

prepared ferrites might be suitable for the application in multilayer chip inductor because

of its good magnetic properties and low loss at high frequency.

Kumar et al. [106] have synthesized magnesium-manganese ferrites having

composition Mg0.9Mn0.1Al0.3CozFe1.7-zO4 (z = 0.3-0.7) by the citrate precursor method.

Single-phase spinel structure of the samples is assured by the X-ray diffraction analysis.

The synthesized materials have been investigated for their electrical and magnetic

properties. Fairly constant value of initial permeability in a wide frequency range of 0.1-

20 MHz along with a low loss factor of the order of 10-4

-10-5

are the superb achievements

of the present study. Moreover, initial permeability increased with an increase in

32

temperature while RLF was found to be low at measurement temperatures. The DC-

resistivity and Curie temperature were observed to increase with dopant content.

Ghatak et al. [107] measured the alternate current conductivity, direct current

conductivity and dielectric properties of Mg-Zn ferrite below room temperature. The

nanocrystalline ferrite powder was prepared from oxides of magnesium, zinc and iron by

using a solid state processing method. The X-ray diffraction analysis was used to

determine the structure and composition of synthesized ferrite, while impedance analyzer,

liquid nitrogen cryostat and electromagnet were used for conduction and dielectric

properties of ferrite. The frequency exponent (s) of conductivity is found to be highly

temperature dependent. The temperature exponent (n) of dielectric permittivity decreases

with increasing frequency. The grain boundaries have more contribution as compared to

the grains contribution in conduction phenomena and resistance due to grain and the

grain boundary contribution possesses two activation regions.

Gabal and Bayoumy [108] synthesized the nanocrystalline Ni0.8-

xZn0.2MgxFe2O4 ferrites (x = 0.0-0.8) using metal nitrates and freshly extracted egg-white.

The synthesized materials were characterized using XRD, FT-IR and TEM. The lattice

constant increased and X-ray density decreased with increasing Mg content. The average

crystallite size determined from XRD data found in the range of 35-59 nm. The

agglomerated particles obtained from TEM analysis are in a good agreement with

crystallite size obtained from XRD. Magnetic properties measured by vibrating sample

magnetometer (VSM) show a decrease in MS up to dopant content of 0.6. The decrease in

the value of coercivity with an increase in dopant content can be explained on the base of

magneto-crystalline anisotropy.

Bharti et al. [109] synthesized magnesium ferrite, zinc ferrite, and magnesium

zinc ferrite using an economic wet chemical synthesis technique. Phase formation of the

synthesized powders has been confirmed by infrared spectroscopy and X-ray Rietveld

refinement analyses. The structural features of these materials have been correlated with

their magnetic properties. Single phase magnesium zinc ferrite nano-particles were

investigated for carbon monoxide and hydrogen gas sensing properties. The response and

recovery transients of conductance were modelled using Langmuir adsorption kinetics

having two active sites in the sensing elements named as CO sensors. The activation

33

energies for response and recovery behavior of these two adsorption sites were found to

be different. This difference in activation energies for response and recovery is due to

different chemi-adsorbed oxygen species in these two sites.

Ghasemi et al. [110] prepared the CuxMg0.5-xZn0.5Fe2O4 (x = 0-0.5) nanoparticles

and thin films by sol-gel route. The morphological features of nanoparticles were

analyzed by TEM. The Mössbauer spectroscopy was dealt with the site preference of the

elements. Magnetic dynamic study was performed by AC magnetic susceptibility

measurements at different frequencies. The phenomenological Néel-Brown and Vogel-

Fulcher models were employed to differentiate the interacting or non-interacting system.

The XRD analysis confirmed the fabrication of single-phase cubic spinel structure. AFM

was used to evaluate the surface morphology of the prepared thin films. VSM was used

for magnetic properties of the samples. The saturation magnetization and initial

permeability were observed to increase with increasing amounts of copper.

Dalt et al. [111] investigated the synthesis of nanostructured MgFe2O4 through

solution combustion technique. The 30% fuel-deficient formulation was selected to

synthesized powders at different furnace temperatures. The structural and morphological

characterizations were performed by using X-ray diffraction and Transmission Electron

Microscopy (TEM). Mössbauer spectroscopy and vibrating sample magnetometer (VSM)

were employed for magnetic measurements. Crystallite sizes of MgFe2O4 around 42.8 nm

calculated from the XRD pattern were consistent with the results obtained from TEM

analysis.

Bachhav et al. [112] synthesized mixed spinel ferrites having general chemical

formula NixMg0.5-xCu0.1Zn0.4Fe2O4 (x = 0.1-0.5) by standard double sintering ceramic

method. The XRD analysis revealed that lattice parameter decreases with increase in Ni

content. They also investigated the temperature dependency of magnetic and electrical

properties of the synthesized ferrites. The variation in initial permeability with frequency

is studied. The Curie temperature (TC) was estimated from initial permeability and AC

susceptibility measurements.

Sadhana et al. [113] have prepared the nanocrystalline MgCuZn-ferrites with

particle size ~30 nm using microwave-hydrothermal method. The structural and

34

morphological features are investigated using X-ray diffraction and scanning electron

microscopy. The grain sizes of the samples are in the range of 60-80 nm. The ultrasonic

velocities measured on MgCuZn-ferrites are found to decrease with an increase of

temperature. The anomaly observed in the thermal variation of longitudinal velocity and

attenuation is explained on the base of magneto-crystalline anisotropy constant.

Mansour [114] prepared the Manganese-magnesium ferrite nanoparticles with

general formula Mn1-xMgxFe2O4; 0≤x≤0.25 using the co-precipitation route. The XRD

analysis confirms the single phase spinel structure. The crystallite size calculated from

Scherrer formula is found to be in the range of 3-6 nm and observed to increase with the

increase in dopant contents. TEM was also utilized to analyze the microstructure of

nanosized Mn1-xMgxFe2O4. Hysteresis loops measured at room temperature revealed the

lower value of saturation magnetization associated with Mn1-xMgxFe2O4 nanoparticles.

Mansour and Elkestawy [115] synthesized the nano-sized Mn1-xMgxFe2O4 (x =

0.0-0.25) by co-precipitation method and reported the effect of Mg-substitution on

structural and dielectric properties of the prepared samples. X-ray diffraction analysis

revealed the nanocrystalline nature of the prepared samples, having crystallite size in the

range of 3-6 nm in nano-samples and 63.9-85.5 nm in bulk samples. The dielectric

properties have been studied in the frequency range of 10-105 Hz at various temperatures.

Dielectric parameters like dielectric constant (έ) dielectric loss (Ɛʺ), dielectric loss tangent

(tan δ) and AC conductivity have been determined for the prepared samples. It is evident

from the obtained data that the magnitude of dielectric constant and loss factor of former

are ten times greater than those of the latter. The low dielectric behavior makes these

materials suitable for high frequency applications.

Khan et al. [116] synthesized Mg1-xTbxFe2O4, ferrite (0.0≤x≤0.2) by ceramic

method and employed X-ray diffraction, Fourier transform infrared spectroscopy (FTIR)

and vibrating sample magnetometer for characterization. The XRD patterns indicate the

single phase formation up to x≤0.04. The lattice constant shows slight increase with the

increase in terbium content up to x = 0.04 and decreases for x > 0.04. This increase

corresponded to the difference in the ionic radii of the cations involved and led to the

formation of secondary phase (TbFeO3). The bulk density increased from 3.5 to 4.6

35

(g/cm3) with the increase of terbium contents. The FTIR spectra showed two significant

absorption bands in the range of wave number 370-1500 cm-1

which assure the spinel

structure and the completion of the chemical reaction. The magnetic properties revealed a

decrease in the saturation magnetization depending upon the Tb content. An unexpected

increase of the saturation magnetization for Tb content of 0.02 could be due to shifting of

Mg ions towards A-sites, in agreement with the results of FTIR. The resistivity increased

with the increase in terbium contents as compared to the undoped sample while the drift

mobility was observed to decrease.

Kaiser [117] studied the structure, electric and dielectric properties of Indium

substituted Mg-Cu-Mn ferrites with a general formula of Mg0.9Cu0.1Mn0.1InxFe1.9-xO4 (x =

0.0-0.4). X-ray diffraction (XRD) analysis indicates the formation of single-phase cubic

spinel structure up to 0.2 and mixed phase (cubic and tetragonal phase) for x ≥ 0.3. The

electrical conductivity of samples revealed a transition from semiconductor to metallic

character with increasing In3+

concentration. Temperature dependence of universal

exponent (S) indicates the presence of two hopping conduction mechanisms: the

correlated barrier hopping (CHB) at low In3+ content x ≤ 0.1 and small-polaron (SP)

hopping at relatively higher In3+

content x ≥ 0.2. The change in the dielectric permittivity

(έ, ɛ) with the temperature at different frequencies shows a normal behavior, whereas the

variation in dielectric loss tangent with frequency at different temperatures shows

unusual behavior with more peak relaxation. The conduction mechanism in the

synthesized ferrites has been discussed in the light of electron exchange between Fe3+ and

Fe2+

ions and hole hopping between Mn3+

and Mn2+

ions at the B-sites.

Shah et al. [118] intended to improve the humidity sensing properties of Mg-

ferrite by doping with Pr in 0.1 mol% and 0.3 mol% concentrations. The spin density

calculated from electron paramagnetic resonance (EPR) increased from 8.15×1020

to

15.6×1020 for 0.3 mol% Pr doped compared to undoped Mg-ferrite. The bulk porosity of

the samples increases from 8.4 to 34% with Pr contents. Pr doping also caused the

increase in sensitivity factor Z10%/Z90% from 24 to 113. Impedance gradient |dlog Z/dRH|

at low 10-30% RH and high 70-90% RH was determined corresponding to spin density

and porosity of the samples. A maximum drift in humidity hysteresis of 22% RH is

36

mainly reduced to 2% RH by a 0.1 mol% Pr doping in magnesium ferrite. Such

significant improvements make this material as a strong candidate for humidity sensors.

Ferreira da Silva and Valente [119] prepared the glasses with composition

xFe2O3-5MgO-(95-x) SiO2, (x = 1.25, 2.5, 5 and 10 mol %) by the sol-gel method. The

precursor was treated at temperature between 500 and 1000 oC. In most of the samples

the presence of magnesium ferrite was detected by XRD. However, the brown and

transparent 1.25Fe2O3-5MgO-93.75SiO2 sample treated at 500 oC was free of magnesium

ferrite. In 10Fe2O3-5MgO-85SiO2 samples treated at 500 and 1000 oC, hematite was

detected. The average crystallite size of the magnesium ferrite treated at 1000 oC, found

in the range of 8-10 nm. All the samples treated at 1000 oC exhibited ferro/ferrimagnetic

interactions combined with superparamagnetism except 1.25Fe2O3-5MgO-93.75SiO2

sample treated at 500 oC which showed a paramagnetic nature down to 5 K.

Patil et al. [120] reported the synthesis of spinel MgFe2O4 by a simple,

inexpensive combustion method and applied as a gas sensor for reducing gases (LPG,

Acetone, Ethanol, Ammonia). The reducing gas sensing properties as a function of

structural and surface morphological properties has been studied. The structural and

morphological features were analyzed by X-ray diffraction and scanning electron

microscopy, respectively. The porous morphology revealed by SEM analysis owed to

decrease with the grain growth by an increase in sintering temperature. The maximum

response of 71% to 2000 ppm of LPG was observed at 698 K with the synthesized

material.

Mounkachi et al. [121] prepared the Mg0.6Cu0.4Fe2O4 ferrites using the solid-state

reaction technique. The structural properties have been studied using X-ray diffraction

analysis. While magnetic measurements were carried out using mean field theory and

high-temperature series expansions (HTSE), extrapolated with the padé approximants

method. The Mössbauer data were dealt to compute nearest neighbor super-exchange

interactions for intra- and inter-site using the probability approach. The obtained

experimental results are in good agreement with the theoretical ones obtained by the

magnetic measurements.

Mohammed et al. [122] synthesized the Mg-Zn mixed ferrites by using standard

ceramic method with general formula, Mg1-xZnxFe2O4 (0≤x≤1). The structural properties

37

of the synthesized materials were elucidated by using XRD and IR spectroscopy. The

various parameters like lattice constant, density, particle size, force constants, bond

length, porosity, shrinkage and cation distribution were estimated. The aggregation of

crystallites of about 200-800 nm is evident from micrographs of scanning electron

microscope (SEM) and transmission electron microscope (TEM). Far infrared absorption

spectra indicates two significant absorption bands and the wave number of the first band,

υ1, decreases while that of the second band, υ2 shifts linearly towards higher wave

numbers with the increase in Zn contents. The room temperature electrical resistivity was

observed to decrease with Zn contents. Data of the vacancy model parameters showed

that the packing factors Pa and Pb decrease, the fulfillment coefficient, α, remain

constant, while the vacancy parameter, β, has significant increase with the increase in Zn

contents. The observed behavior of the above-said parameters indicate the presence of

cation or anion vacancies and the partial contribution of the Zn2+ vacancies in the

enhancement of the electrical conductivity of Mg-Zn ferrites.

Kaur et al. [123] used the oxalyl dihydrazide-metal nitrate combustion method to

synthesize the Mg1-xCdxFe2O4 (x = 0.0-0.6) nanoparticles (NPs). Synthesized NPs were

subjected for characterization by various techniques namely XRD, SEM and TEM.

Magnetic measurements were performed by using vibrating sample magnetometer

(VSM). The synthesis method used is revealed to be a low temperature route for the

preparation of mono-dispersed ferrite NPs with average particle size in the range of 22-34

nm. The magnetic parameters like saturation magnetization and remanence increased

with cadmium content up to x = 0.4 (Mg0.6Cd0.4Fe2O4) and these ferrites could be a

potential soft magnetic material with promising magnetic and micro-structural properties.

The temperature of phase formation (300 oC) is much lower than the earlier reported

value (700 oC) for co-precipitation method. In addition, the combustion approach does

not require sintering and milling process at high temperature (as required by the

conventional ceramic method) which causes lattice defects, strains and the magnification

of the ferrites.

Khan et al. [124] studied the effects of terbium (Tb) doping on several properties

including ferromagnetic resonance (FMR), relative initial permeability, morphology, and

dielectric properties of MgFe2O4. Tb doping results in the fluctuation of FMR line-width

38

and FMR position of magnesium-terbium (Mg-Tb) ferrites. The relative initial

permeability showed a progressive decrease with increasing Tb content. The magnetic

loss factors of substituted samples show a decreasing trend with the increase in frequency

from 1 kHz to 10 MHz. The doped samples have a lower dielectric constant with respect

to the pure MgFe2O4 ferrite. The lower dielectric constant is attributed to the less ease of

electron exchange mechanism caused by lockup between iron and terbium ions. The

observed irregular behavior in dielectric loss is related to the hopping probability of

charges in synthesized materials. Ac conductivity is observed to decrease beyond a

certain frequency of 8 MHz which attributed to the occurrence of dielectric loss.

Singh et al. [125] investigated the DC and Ac electrical resistivity of Mg

substituted Ni-Cu ferrites as a function of temperature, frequency and composition. The

Ac-resistivity of the samples decreases with an increase in frequency, having

ferrimagnetic behavior. The dielectric loss tangent showed maximum frequency

dependence in between 10Hz and 1 kHz in all the samples. The conductivity relaxation of

charge carriers has been examined using the electrical module formulism, and results

indicate the presence of the non-Debye relaxation in the synthesized samples. Similar

values of Ea both for DC conduction and conductivity relaxation indicate that the

mechanism of electrical conduction and dielectric polarization is the identical in these

ferrites. The saturation magnetization and coercivity calculated from the hysteresis loops

show striking dependence on composition.

Sujatha et al. [126] examined the effect of Mg substitution on structural and

magnetic properties of Ni-Cu-Zn ferrites prepared by sol-gel process. Various parameters

were studied using X-ray Diffraction (XRD), Fourier transform infrared spectroscopy

(FT-IR), Field Emission Scanning Electron Microscopy (FE-SEM) and Vibrating sample

magnetometer (VSM). XRD analysis was used for the phase identification, unit cell

parameter and crystallite size determination. The lattice constant is found to decrease

with the increase in Mg content. Saturation magnetization and coercivity showed an

opposite behavior with the Mg content. The Curie temperature (TC) increases while initial

permeability (µi) decreases with the increase in Mg contents. This is because of reduced

magnetization, grain size and enhanced magneto-crystalline anisotropy. The permeability

is also found to be constant up to 30MHz frequency ensuring the compositional stability

39

and quality of the material. The synthesized materials have attraction for applications in

Multilayer Chip Inductors because of their invariable permeability and high thermal

stability.

Albuquerque et al. [127] studied the structural and catalytic properties of Co-Cu-

Ni ferrites. Nanostructured ferrites were synthesized by co-precipitation method with

particle size in the range of 3-10 nm. The chemical and structural characterizations were

performed using X-ray diffraction, X-ray fluorescence, X-ray photoelectron spectroscopy

and Mössbauer spectroscopy. The catalytic efficiency of the samples was tested by the

decomposition of hydrogen peroxide and the oxidation of methylene blue, monitored by

UV-vis spectro-photometry. It was observed that the presence of cobalt ions is a crucial

factor to achieve a systematic effectiveness of the catalyst in the decomposition of H2O2.

On the other hand, Cu ferrites gave a better performance in the oxidation of methylene

blue, attributed to the different redox properties of Cu with ease of electrons to contribute

in the oxidation of organic compounds.

Chen et al. [128] synthesized the Mg-ferrite nanoparticles with high saturation

magnetization using microwave assisted ball milling process. The as-milled materials

were characterized by X-ray diffraction, transmission electron microscopy and vibrating

sample magnetometer. The obtained results revealed that the average size and the

saturation magnetization of Mg-ferrite nanoparticles were 30 nm and 43.40 emu/g,

respectively. The microwave assisted ball milling is a simple and environment‐friendly

approach as compare to conventional milling process and could be considered a

promising approach for the synthesis of nanoparticles of high performance in the future.

Bobade et al. [129] synthesized the nanocrystalline Ni2+ substituted Mg-Zn

ferrites with general formula Mg0.7-xNixZn0.3Fe2O4 (x = 0.0-0.6) by using sol-gel auto-

combustion method. The citric acid was used as a fuel and the metal nitrate to citric acid

ratio was taken as 1:3. The structural and morphological features of Mg-Ni-Zn ferrites

were studied by X-ray diffraction, scanning electron microscopy, and FT-IR

spectroscopy. The lattice parameters were determined from the XRD data. The FTIR

spectroscopy is used to deduce the cation distribution between tetrahedral and octahedral

sites of spinel Mg-Ni-Zn ferrites. Micrographs indicate the grain growth formation with

40

the increase in Ni2+

contents. The M-H loops were utilized to find out the saturation

magnetization (MS) and magneton number of the synthesized materials. The value of Ms

increases with the increase in Ni2+ contents of Mg-Zn ferrite.

Sujatha et al. [130] prepared the nanoparticles of Mg0.5Cu0.05Zn0.45Fe2O4 through

sol-gel method using polyvinyl alcohol as a chelating agent. The as-synthesized sample

was annealed at three different temperatures (500 oC, 700 oC and 900 oC). The phase

formation, morphology and magnetic properties were studied using X-ray diffraction

(XRD), Fourier transform infrared spectroscopy (FTIR), field emission scanning electron

microscopy (FESEM) and vibrating sample magnetometer (VSM). The crystallite size

and magnetization increased with annealing temperature. The coercivity increased up to a

specific annealing temperature and decreased afterwards, showing transition from single

domain to multi domain state with increasing annealing temperature. In addition, to

assess the capability of the material as a ferrite core, in multilayer chip inductors, the

powder sample annealed at 500 oC was compacted in shape of torroids and sintered at

three different temperatures (800 oC, 900

oC and 950

oC). The permeability increased

with an increase in sintering temperature indicating dependency on microstructure. The

frequency dispersion of the permeability, for the sintered samples, demonstrated higher

frequency stability. The cut-off frequency for the samples sintered at 800 oC, 900

oC and

950 oC is 32 MHz, 30.8 MHz and 30.4 MHz, respectively.

In recent years, nanomaterials have emerged as a rapidly advancing field,

providing vast avenues of research. Rapid growth in the demand for telecommunication

and high-frequency magnetic devices require production of materials with significant

improvements in their performance with lower fabrication costs. The literature of the last

10-15 years reveals that spinel magnesium ferrites of different compositions have been

extensively studied and used in technological products. However, the search for a better

product with lowest energy consumption and optimum performance is still going on.

The various properties of spinel Mg-ferrite systems have been investigated by

doping with single metal e.g. Cr, Al, Gd, Cu as well as co-substitution of two different

metals e.g. Cu-Zn, Gd-Co, Nd-Co and Zr-M (M = Mn and Zn). The electrical and

magnetic properties were found convincing for various electromagnetic applications.

Verwey et al., found the relationship between the electronic conduction and distribution

41

of cations at different sites. It was found that the electronic exchange taking place

between ions of the same element with different valence state at octahedral sites in the

crystal lattice called hopping conduction mechanism. The DC-electrical resistivity of Mg-

ferrite is reported as 104-10

7 Ωcm which decreases as the temperature increases;

semiconductors like behavior.

On the base of literature survey carried out for magnesium ferrite spinel; it is

observed that the efforts of the researchers have been mainly concentrated on the

synthesis of single crystals, bulk materials, thin films and nanosized structures by using a

variety of well known methods including the coprecipitation method, sol-gel and solid

state method. The physical, magnetic and electrical properties of ferrite materials depend

upon the synthesis routes, chemical composition, annealing temperature and distribution

of metal ions at A- and B-sites. The conventional ceramic method of preparation involves

the mixing of suitable metallic oxides with appropriate grinding followed by a solid-state

reaction at high annealing temperatures of 1573-1973 K. Though the route is quite simple

yet it has several drawbacks, such as; high sintering temperature, large reaction time,

large particle size and limited degree of homogeneity. However, chemical methods like

co-precipitation, spray-pyrolysis and sol-gel etc., are preferable because apart from the

advantages of an economical and low-temperature processing, these methods make it

possible to obtain nanosized ferrite materials. Although the aforesaid methods are capable

of producing nanometric ferrites, but the quality of producing nanoparticles in many

cases has a large size distribution along with the arbitrary size control. In most of these

methods, size variation is achieved through annealing at various temperatures. If multiple

sizes are produced by the synthesis without thermal annealing or without changing the

thermal annealing temperature, often major changes in the synthesis procedure is required

[131]. To correlate size effects with changes in magnetic and electric properties, it is

essential to select a synthesis method which allows controlling the size of nanoparticles

with a narrow size distribution [132]. Microemulsion method is preferred over other

methods for the reason stated above and has many advantages; especially low

temperature wet chemical synthesis to obtain single-phase materials. It is also easier and

cost effective than the other chemical methods. The microemulsion method has been

42

known to successfully prepare nanocrystalline spinel ferrites with the particle sizes less

than 50 nm.

1.8 Objectives and Plan of Work

The aims and objectives of the present investigation were:

• To synthesize magnesium ferrite with a size < 50 nm in order that the signal to

noise ratio could be reduced for use in recording devices [133].

• To prepare ferrite materials able to (i) sustain stability in different

environments and at high temperatures and (ii) having high Curie

Temperature (TC)

• The pre-requisite for ferrite materials useful in magnetic devices is large

saturation magnetization along with a moderate coercive field. One of the

objectives of this research work is to enhance the saturation and remanence

magnetization of Mg-ferrite by doping with different metal cations.

• Technical requirements of the materials suitable for applications in microwave

devices are high values of electrical resistivity (>107 Ω.cm) to curb the eddy

current losses and low values of dielectric loss. Thus, aim of the present

research work is to enhance the electrical resistivity and to reduce the

dielectric constant and dielectric loss of doped Mg-ferrites to make it suitable

and useful in microwave devices.

After going through the relevant literature on the subject, it becomes evident that

less attention is given to the modification of magnetic and electrical properties of doped

Mg-ferrites. In order to achieve the above-mentioned objectives of the research work,

substitution of Mg-ferrite (MgFe2O4) matrix with binary mixtures of transition metals

ions, such as Co-Cr, Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr at the Mg2+

and Fe3+

sites. The

following five series of the doped Mg-ferrites are synthesized; Mg1-xCoxCrxFe2-xO4, Mg1-

xNixCrxFe2-xO4, Mg1-xCuxCrxFe2-xO4, Mg1-xZnxCrxFe2-xO4 and Mg1-xMnxCrxFe2-xO4 (x =

0.0-0.5).

43

The selection of the dopant metals i.e. M-Cr (M = Co2+

, Ni2+

, Cu2+

, Zn2+

and

Mn2+

) is based on their several features like ionic radii, number of unpaired electrons and

significantly higher electrical resistivity etc. Since Cr3+ ions are less conducting with high

electrical resistivity so that their substitution in Mg-ferrite is expected to enhance the

electrical resistivity and reduce the dielectric loss of the synthesized materials. Moreover,

Cr3+

being a transition metal would preferentially occupy octahedral lattice sites [134], so

that, when substituted for iron it could enhance the electrical resistivity by decreasing the

number of iron ions at the octahedral sites. The resistivity enhancement is a challenging

task for researchers working in the field of ferrites. High resistive materials are

beneficial for applications in microwave devices [135]. Few reports are available on the

electrical and dielectric properties of spinel Mg-ferrite nanomaterials. The study of

dielectric parameters and DC-electrical resistivity provide valuable information on the

behavior of localized and free electric charge carriers. This leads to a better

understanding of the conduction mechanism.

The substitution of transition metal (M = Co2+

, Ni2+

, Cu2+

, Zn2+

and Mn2+

) ions

with specific site occupancy in Mg-ferrite could tune saturation magnetization,

remanence magnetization, hyperfine magnetic interaction, cubic anisotropy and

coercivity. Therefore, doping by binary mixtures of M-Cr is expected to attain the

objectives of present research. The variation of hyperfine parameters (center shift,

quadrupole splitting and hyperfine magnetic field) can further confirm the presence of

doped cations in the Mg-ferrite lattice at different crystallographic lattice sites. Ionic radii

of Co2+

, Ni2+

, Cu2+

, Zn2+

, Mn2+

and Mg2+

are 0.72, 0.69, 0.73, 0.82, 0.67 and 0.72Å,

respectively, and those of Cr+2

and Fe3+

are 0.63 Å and 0.64 Å, respectively. Therefore,

these ions having comparable ionic sizes would replace the host cations easily resulting

in the formation of novel doped compounds without any phase distortion of the cubic

spinel lattice.

The correlation of magnetic and electrical properties with dopant contents would

be established by an investigation of various compositions of substituted Mg-ferrites.

Moreover, the effect of cationic substitution on the various parameters namely center

shift, quadrupole splitting, hyperfine magnetic field, magnetization, first-order cubic

44

anisotropy, coercivity, Curie temperature, electrical resistivity, dielectric dispersion and

dielectric loss, would be investigated. The study reported here is perhaps the first of its

kind that attempts to address the Mg-ferrite (MgFe2O4) and its allied compounds (Mg1-

xMxCrxFe2-xO4) in nano regime with different cationic substitutions.

CHAPTER 2

EXPERIMENTAL

44

2. EXPERIMENTAL

2.1 Methods of Preparation

Novel characteristics and numerous applications of nanoscale materials in various

scientific fields such as physical, chemical, biological and engineering sciences demand

the development of new synthetic routes. In the last few years, many publications have

described efficient synthetic routes to shape-controlled, highly stable, and mono-

dispersed magnetic nanoparticles. Several popular synthetic routes have been fabricated

for the synthesis of spinel ferrite nanomaterials including sol-gel, coprecipitation,

hydrothermal, spray-pyrolysis, sonochemical and microemulsion methods which are

described as follows.

Sol-gel method [136, 137] refers to the hydrolysis and condensation of metal

alkoxides-based precursor. Metal alkoxides are dissolved in a suitable solvent to obtain a

sol and loss of solvent leads to the transition from the liquid sol into a solid gel

phase. This wet gel is converted into a dense and dry precursor with further drying and

heat treatment. Required product is then obtained after annealing the dried precursor at

high temperature. The size of the sol particles depends on the control of several factors,

such as solution composition, pH and temperature [138]. Several chelating agents (citric

acid, stearic acid, poly vinyl alcohol (PVA) and ethylene glycol) are also used by

different researchers in the synthesis process [139-146]. The disadvantage of this method

is long preparation time.

Co-precipitation method [147-149] involves nucleation, growth and

agglomeration processes simultaneously. When precipitation begins, numerous small

crystallites are formed (nucleation) that tend to aggregate rapidly to form larger particles,

which is not a good practice in the synthesis of nanomaterials. The synthetic reaction of

the oxide can be generally divided into two categories: (i) first which can produce an

oxide directly; (ii) that which produces a precursor that must undergo further processing

(drying, calcinations). Nanosized ferrites are usually prepared in an aqueous medium

whose chemical reaction of formation may be written as follows:

M2+

+ 2Fe3+

+ 8OH¯ → MFe2O4 + 4H2O

Where M can be Mn2+

, Co2+

, Cu2+

, Mg2+

, Zn2+

and Ni2+

. Complete precipitation in

the form of metal hydroxides should be expected at a pH level between 8 and 14 by the

45

addition of a precipitating agent such as NaOH, NH4OH or a solution of Na2CO3, with a

stoichiometric ratio of 2:1 (Fe3+

/M2+

) in a non-oxidizing oxygen environment [150]. The

resulting hydroxides are annealed after filtration and washing to get the final oxide

powder [151]. The size, shape, and composition of the magnetic nanoparticles very much

depend on the type of salts used (e.g. Chlorides, sulfates, nitrates), the Fe2+

/Fe3+

ratio, the

reaction temperature, the pH value and Ionic strength of the media.

Hydrothermal method [152-154] based on the principle that in a sealed vessel, i.e.

bomb or autoclave, the solvent (water) is brought to temperature above their boiling point

by increasing the pressure as a result of warming. When a solvent is handled in such

conditions, it is designated as a solvothermal process. The reaction between metal salts

solution can be carried out in water or any suitable solvent. In case of water, the process

is referred as hydrothermal process. The critical point of water is at 647 K and 219

atmospheric pressure. Above this temperature and pressure, water is said to be a

supercritical and possess characteristics of both liquid as well as gas. Many inorganic

materials such as TiO2 [155], BaTiO2 [156], CoFe2O4 [157], ZnFe2O4 [158] and

SrFe12O19 [159] have been synthesized by this method. The disadvantage of this method

is slow reaction rate at any given temperature. In order to increase the crystallization

kinetics, it is necessary to introduce the microwave, electric or ultrasonic fields in the

hydrothermal systems and are referred to as microwave-hydrothermal, electrochemical-

hydrothermal and ultrasonic-hydrothermal, respectively [160-162] which contribute to

make this method expensive and time consuming.

The spray-pyrolysis method [163, 164] entails the production of aerosol droplets

by the atomization of the starting metal salts solution, sol or suspension. The droplets

generated undergo evaporation with solute condensation within the droplet and then

performed drying, thermolysis of the precipitated particles at high temperatures to form

micro-porous particles, and finally sintering to form dense particles. This is a useful

method for the synthesis of high purity nanoparticles homogeneously. The main

disadvantage of spray pyrolysis is the use of large amounts of solvent resulting in

significant production cost and difficulty in the intensification of product.

Sonochemical method [165] was originally proposed for the synthesis of iron

nanoparticles, but today it is commonly used to synthesize nanoparticles of different

46

metal oxides [166] with the aid of controlling the coating depth of target materials by

adjusting reaction conditions such as the amount of reactants and/or the sonication time

[167, 168]. Various theories have been established to explain how 20 KHz sonic radiation

can break chemical bonds. However, all these theories agreed that, when liquids are

irradiated with ultrasonic irradiation, acoustic cavitation will form. Ultrasonic cavitation

is the formation, growth, and implosive collapse of bubbles in a liquid. Ultrasonic

cavitation produces a variety of physical and chemical effects, such as high temperature

(>5000 K), pressure (>20 MPa), and cooling rate (>1010 Ks-1), which could provide a

unique environment for chemical reactions under extreme conditions. A very high

temperature (5000-25000 K) obtained by collapsing bubbles of liquid is responsible for

the breaking of chemical bonds [169]. In most of cases, the end product is amorphous and

highly aggregated so it is very difficult to characterize individual particles [170]. The

main disadvantage is that it requires expensive equipment.

Although, the methods described above are capable of producing nanoscale

ferrites, but the quality of nanoparticles is often poor in many cases due to a large size

distribution along with the arbitrary size control. In most of these methods, size variation

is achieved through annealing at various temperatures. But it has been observed that the

cation distribution at A- and B-sites is strongly affected by the annealing temperature

[171]. Due to ferrimagnetic nature, the magnetic properties of spinel ferrites depend

strongly on the distribution of cations between the A-and B-sites. When the annealing

temperature is used to control the size of nanoparticles, a direct correlation between the

size effect and the magnetic response is not possible because of the additional variable of

cation redistribution. Although multiple sizes could be produced by the synthetic process

without thermal annealing or without changing the thermal annealing temperature, often

major changes in the synthesis procedure are required [131].

To correlate size effects with changes in magnetic and electric properties, it is

essential to select a synthesis method which allows control over the size of nanoparticles

with a narrow size distribution [132]. Microemulsion method [172, 173] has many

advantages; especially low temperature wet chemical synthesis to obtain single-phase

materials. It is easier and cost effective over other chemical methods. The product yield is

high compared to the other methods described above and also has the advantages of

47

reliability and reproducibility. Microemulsion procedures are commonly used to build

high-quality metal and semiconductor nanoparticles with narrow size distribution [133,

174-176]. Moreover, by minor adjustments to the conditions of synthesis, size control can

be achieved easily. In the microemulsion method, surfactant concentrations above the

critical micelle concentration (CMC) are usually used. In the present work, polyethylene

glycol assisted microemulsion method [177] has been adopted for the synthesis of Mg-

ferrite and its metal doped derivatives.

2.1.1 Microemulsion method

Synthesis of nanoparticles by microemulsion method [146] involves surfactant, a

combination of two chemicals i.e. one of them is soluble in water and the other in the

organic phase only. Hoar and Schulman [178] observed that certain combinations of

water, oil, surfactant and alcohol or the amine-based co-surfactant make clear, apparently

homogeneous solution which is known as microemulsion.

An emulsion is made by mixing a small amount of water in a large volume of the

organic phase and a surfactant is added. The size of water droplets is directly related to

the proportion of the water to surfactant. Surfactant molecules accumulate on the surface

of the water drop and stabilize the drop. Such a droplet is termed as a reverse micelle.

Since the drop is small, only a small amount of reagents can be squeezed into it. When

this droplet reacts with another reactant, a tiny particle is formed. The domain size of the

dispersed phase in microemulsion is usually very small (a few nanometers) and chemical

reactions can occur within the nano-droplets of water or at the oil-water interface in the

microemulsion [179].

The special physicochemical properties motivate the material scientists to use

microemulsion for the synthesis of new materials, nano-sized particles in particular [180].

The selection of surfactant used in the formation of microemulsion for the synthesis of

nanomaterials is very important. The selection of a surfactant is based on several factors

including reactants and reaction conditions. The water-in-oil microemulsion technique

has been used for the synthesis of various nanoparticles, including metals [181-183],

halides [184] and oxides [185, 186]. A variety of surfactants such as

cetyltrimethylammonium chloride (CTAC), sodium dodecyl sulfate (SDS) and

48

polyvinylpyrrolidone (PVP) have been used by different researchers for the synthesis of

nano-ferrites [138, 146, 187].

2.2 Chemicals Used

The chemicals used for the synthesis of the samples were of high purity procured

from well known suppliers. Therefore no further purification was considered necessary

and hence used as received. The details of the chemicals are as follows:

Table 2.1: Specifications for the chemicals used

S. No. Compound Chemical Formula Purity Supplier

1 Ammonium hydroxide NH4OH ≥97 % Fluka

2 Chromium nitrate Cr(NO3)3 99.0 % Reidal

3 Cobalt acetate Co(CH3COO)2 99.0 % Fluka

4 Copper acetate monohydrate Cu(CH3COO)2·H2O 99.0 % Merck

5 Ferric nitrate nonahydrate Fe(NO3)3.9H2O 99.0 % Merck

6 Magnesium nitrate hexahydrate Mg(NO3)2.6H2O 99.9 % Merck

7 Manganese acetate tetrahydrate Mn(CH3COO)2.4H2O 98.0 % Merck

8 Nickel acetate tetrahydrate Ni(CH3COO)2.4H2O 99.0 % Merck

9 Polyethylene glycol PEG 98.0 % BDH

10 Zinc acetate tetrahydrate Zn(CH3COO)2.4H2O 99.0 % Merck

2.2.1 Samples preparation procedure

The spinel magnesium ferrite having nominal composition MgFe2O4 is prepared

by a polyethylene glycol (PEG) assisted microemulsion method [177]. The desired

metallic salts i.e. Mg(NO3)2.6H2O and Fe(NO3)3.9H2O, are dissolved in de-ionized water

and mix together in an appropriate stoichiometric ratio in a beaker. Aqueous solution of

polyethylene glycol (PEG) has been added in 1:2 ratio of the mixed metal salt solution:

PEG solution. Ammonium hydroxide solution (0.2 M) is added for the precipitation at pH

9-10. The mixture is stirred for 3 hrs with a slight warming followed by an overnight

aging. The collected precipitates comprising of hydroxides of metal ions,

nitrates and some water contents

water and ethanol in order to remove the

species, after that dried at 40

calcination at 573 K for 2 hrs

in a set temperature tube furnace

may occur:

Mg(NO3)2 (aq) + 2Fe(NO

yNH4NO3

The grinding of the

by using an Agate mortar

characterizations. The pow

mm in diameter and

electrical resistivity

polyethylene glycol (PEG) assisted microemulsion method

Figure 2.1: Flow sheet diagram

Drying 403 K

The collected precipitates comprising of hydroxides of metal ions,

and some water contents [188], have been repeatedly washed with de

in order to remove the excess ammonia, PEG and

after that dried at 403 K in an oven to remove the water contents

calcination at 573 K for 2 hrs. The dried product obtained is annealed at 1123K for

tube furnace at a heating rate of 5 Kmin-1. The following reaction

(aq) + 2Fe(NO3)3 (aq) + yNH4OH [Mg2+

Fe3+

.OH

[Mg2+

Fe3+

.OH-]m + yNH4NO3

rinding of the obtained crystalline samples into fine powder

by using an Agate mortar-pestle. The obtained material is subjected

The powdered samples are also pressed at 7 ton

diameter and ~2 mm thickness and have been used in measurements

and dielectric properties. The flow sheet diagram for the

polyethylene glycol (PEG) assisted microemulsion method is shown in

Flow sheet diagram for the synthetic scheme.

Drying 403 K

(Crystalline)

Calcination 573 K

49

The collected precipitates comprising of hydroxides of metal ions, ammonium

repeatedly washed with de-ionized

excess ammonia, PEG and any un-reacted

3 K in an oven to remove the water contents, followed by its

s annealed at 1123K for 8 hrs

The following reaction

.OH-]m.nH2O +

[Mg2+

Fe3+

.OH-]m

MgFe2O4

samples into fine powder has been done

s subjected to various

7 ton to form pellets of 13

used in measurements of DC-

The flow sheet diagram for the

is shown in Fig. 2.1.

Annealing 1123 K

(Crystalline)

Precursor Calcination 573 K

50

All series of Mg-ferrites doped with M-Cr (M = Co, Ni, Cu, Zn and Mn) have

been synthesized using the same method as described on previous page, with the addition

of the aqueous salt solution of appropriate molarities of the dopants. The pathway

adopted for the synthesis of substituted Mg-ferrites is shown in Fig. 2.1.

2.3 Characterization Techniques

The following characterizations have been potentially performed for the analysis

of the synthesized samples.

1. Thermo-gravimetric analysis (TGA/DTG)

• Thermal events and phase transformation

2. Powder X-ray diffraction technique (XRD)

• Phase identification

• Lattice parameters determination

• X-ray and bulk density calculation

3. Scanning electron microscopy (SEM)

• Surface morphological and microstructural analysis

4. Energy dispersive X-ray fluorescence (ED-XRF)

• Elemental analysis of the prepared samples

5. Mössbauer analysis

• Center shift

• Quadrupole splitting

• Hyperfine interactions

• Site population

6. Super conducting quantum interference device (SQUID) magnetometer

• Saturation magnetization

• Remanence

• Magnetocrystalline anisotropy coefficient

• Coercive field

7. Vibrating sample magnetometer (VSM)

51

• Curie temperature

8. DC-electrical resistivity measurements by two probe method

• Electrical resistivity at potential operational range around 300 K

• Activation energy

• Drift mobility

9. Dielectric measurements by LCR meter bridge

• Dielectric constant

• Dielectric loss tangent

2.3.1 Thermo-gravimetric analysis

Thermal analysis is based on the determination of changes in sample weight in

relation to changes in temperature in a controlled way [189]. In thermo-gravimetric

analysis (TGA), the mass of a sample in a controlled atmosphere is recorded continuously

as a function of temperature or time with the increase in sample temperature. Two types

of thermal curves are obtained simultaneously in this thermal analysis namely thermo-

gravimetric analysis (TGA) and differential thermo-gravimetric analysis (DTG). Such

thermal (TGA/DTG) analyses rely on a high degree of precision in measurement of

weight, temperature and temperature change. Thermo-gravimetric analysis (TGA)

measures the weight changes in a sample as a function of temperature and this technique

is primarily used for the determination of phase transformation occurring in the materials.

In general, the TGA curves are plotted with the percent weight change against

temperature. A differential weight loss curve (DTG) might be a good protocol that can

identify the point where weight loss is most apparent. A schematic representation of the

operating principle of a thermo-gravimetric analyzer is given in Fig. 2.2.

The analyzer consists of an ultra-micro thermo-balance equipped with digital

balance systems, with top-loading sample arrangement and direct temperature

measurement at the sample. The sample holder (crucible) is placed in a small electrical

oven with a thermometer to sense the temperature. This part has vacuum tight

construction and sophisticated gas paths for inert atmosphere to prevent

oxidation/undesired reactions. The thermo-balance sends the weight loss signal to the

computer for storage, along with the sample temperature and the elapsed time. The TGA

52

signals are used to plot the TGA curve, converted to percent weight change on the Y-axis

against the reference material temperature on the X-axis as shown in the inset of Fig. 2.2.

Figure 2.2: A block diagram of a thermo-gravimetric analyzer [189]

In the present study, thermo-gravimetric/differential thermo-gravimetric

(TGA/DTG) analyses of un-annealed samples have been performed simultaneously at a

default heating rate of 5 Kmin-1 using TGA (Perkin Elmer instrument) in the temperature

range of 323-1373 K.

2.3.2 Structural analysis

Structural analysis was performed in terms of phase identification, morphological

studies and chemical composition of the samples. X-Ray diffraction (XRD) analysis was

carried out using PANalytical 3040/60 X` Pert PRO diffractometer with Cu Kα as a

radiation source at 45 kV and 40 mA. The XRD peaks of the powdered samples were

recorded between 10o and 80

o with a scan step of 0.04

o and step time of 1 s/step. SEM

(Hitachi S-3400 N) was used to observe the microstructure and morphological features of

the synthesized samples. The quantitative elemental analysis was carried out by means of

energy dispersive X-Ray fluorescence spectrometer (Horiba, MESA-500).

53

2.3.2.1 Powder X-ray diffraction (XRD)

Powder XRD is widely used to identify the crystal structure of a substance that

gives the fingerprint image of the crystal structure of this substance. This technique is

mainly used for the identification of phase formation and crystal structure of the different

materials, which is done by matching the obtained peak pattern with that of the standard

pattern of the same substance. In addition, we get information about the structural

parameters such as lattice constant, cell volume, crystallite size and X-ray density of the

desired phase [190]. There are three different methods for the determination of crystal

structure by X-ray diffraction and all three based on Bragg’s law. These are Laue,

rotating-crystal and powder methods. The powder method [191] can be performed by two

different ways namely Debye-Scherrer camera and diffractometer method [192].

The interaction of X-rays with a crystalline sample can be explained by the

Bragg’s law (Eq. 2.1), which gives a relationship between the diffraction angle (Bragg

angle), X-ray wavelength, and inter-planar spacing. According to Bragg, X-rays strike

with a powder sample, a series of parallel planes of crystalline sample act like mirrors

that “reflect” the X-ray beams as shown in Fig. 2.3. The path differences between a pair

of incident rays traveled through the parallel planes are determined by the inter-planar

spacing. As the total path difference is equal to nλ (where n is the order of interference, λ

is the wavelength of the incident X-ray beam), the constructive interference will occur

and a group of diffraction peaks can be observed and give rise to X-ray patterns.

Constructive interference occurs only between those diffracted rays which obey the

Bragg’s law. The Bragg’s law can be expressed as:

= 2 (2.1)

Where, d is the inter-planar spacing for a given set of hkl, λ is the wavelength, and θ is

the Bragg angle shown in Fig. 2.3.

54

Figure 2.3: Illustration of crystal planes and Bragg’s law [191].

Langfordy and Louër [193] reported the comprehensive review in the literature on

the use of powder X-ray diffraction for characterization of materials. Powdered samples

annealed at 1123 K were characterized by X-ray diffractometer. The identification of

spinel cubic phase of all samples synthesized here was conducted by matching the peak

positions and intensities in the experimental diffraction patterns to those in the ICSD

standard pattern database. Several parameters such as crystallite size, lattice constant, cell

volume and X-ray density were calculated from XRD data. The crystallite size of

nanomaterials was derived from the full width at half maximum (FWHM) values of

indexed peaks in the X-ray patterns and calculated using the Scherrer equation [177,

191]:

= (2.2)

Where, λ is the X-ray wavelength and is equal to 1.542 Å, θ is the Bragg angle, and β is

the line broadening at half the maximum intensity (FWHM) in radians. K is shape factor,

equal to ~0.9 for nanocrystallites [191, 194, 195]. The lattice constant (a) and cell volume

(Vcell) were calculated from the following equations:

= ℎ + + (2.3)

! = " (2.4)

The X-ray density (dx) and bulk density (db) were also calculated using following

formulae [34].

55

# = $%&'()*+,, (2.5)

- = ./0 (2.6)

Where, Z is the number of formula units in a unit cell which is 8 for spinel ferrite, Mw the

molecular mass of the sample, NA the Avogadro’s number, m is the mass of the pellet, h

is the the thickness (∼2 mm) and r is the radius (6.5 mm) of the disk shape pellets.

2.3.2.2 Scanning electron microscopy (SEM)

The scanning electron microscopy (SEM) is a qualitative technique to provide

information on the morphology, topographical features, compositional mapping, phase

distribution, crystal orientation and defects of the powders/solid pieces of the samples.

Incident beam of high-energy electrons can be focused on the sample surface using

electromagnetic lenses and it provides a variety of signals that can be recorded by the

detector for the real-space images of materials with resolutions of about a few tenths to a

few nanometers and to simultaneously obtain diffraction information for specific regions

in the images (e.g., small precipitates) as small as 1 nm [196].

The SEM technique involves some specifications for the sample from about 1 cm

to 5 microns in width can be imaged with a magnification ranging from 20X to about

30,000 X, the spatial resolution of 50 to 100 nm. The SEM is also capable of carrying out

analysis of locations of selected points on the sample.

In the present study SEM analysis was carried out by using the sintered pellets

(dimensions as mentioned in Sec. 2.2.1) of the samples. SEM analysis is considered

"non-destructive", i.e. X-rays produced by electron interactions do not lead to loss of

sample volume, making it possible to analyze the same sample repeatedly.

2.3.2.3 Energy dispersive X-ray fluorescence (ED-XRF)

Energy dispersive X-ray fluorescence technique is a multi-elemental and non-

destructive analysis with high-speed qualitative and quantitative analysis in a wide

concentration range.

56

ED-XRF analyzer is based on the investigation of the elements present in the

sample, when these are excited by the electromagnetic radiation interacting with the

sample. The incident beam can excite an electron from an inner shell of an element,

creating an electron-hole which is then filled up by an electron from a higher energy

outer shell, thereby releases energy in form of X-rays. As the X-ray energy is

characteristic of the energy difference between the two shells, and the atomic structure of

the specific element from which these were emitted, this allows the elemental

composition analysis of the many elements in the sample, simultaneously [197].

The block diagram of energy dispersive X-ray fluorescence spectrometer (ED-

XRF) is given in Fig. 2.4. It consists of an X-ray source of radiations, an energy

dispersive detector, an analyzer and a computer. The analyzer has an ability to analyze

the components of solid, liquid and powder at the rate of high speed without damaging

the sample. The analyzer unit consists of a power supply for X-ray tube and a high

precision current and voltage control circuit for the power supply which gives the X-rays

and is controlled by X-ray controller before sending to sample.

Figure 2.4: Block diagram of energy dispersive X-ray fluorescence (ED-XRF)

The characteristics X-rays from the source fall on the sample and then enter to Si

(Li) detector which is connected to the data processing units. An energy dispersive

detector in combination with a multi-channel analyzer is used to simultaneously collect

the fluorescence radiation emitted by the sample and separate the different energies of the

characteristic radiation from each of the elements of the different sample. The detector

acts as a transducer that converts the X-rays photon to charge. The produced charge is

X-ray Tube

X-ray

controller

Sample

Detector Analyzer

Data processing

unit

57

directly proportional to the energy of the X-rays entering the detector. The data

processing unit comprises a high rate pulse processing circuit which measures the signals

from the detector [197] and the software provides easy access to spectrum analysis,

automatic qualitative and quantitative analysis of the sample.

2.3.3 Mössbauer spectrometry

Mössbauer spectrum covers the magnetic interactions that occur in

ferro/ferrimagnetic materials. Mössbauer absorption spectra of ferrites can be recorded

both at normal and low temperature in transmission geometry using a γ-ray source

57Co/Rh. Due to the high energy and width of extremely narrow lines of gamma rays,

Mössbauer spectroscopy is one of the most sensitive technique in terms of energy

resolution, capable of detecting the change in just a few parts per 1011

.

The phenomenon of recoilless emission and resonant absorption of nuclear γ-rays

in solids is known as the Mössbauer effect and was discovered by the German Physicist

Rudolf L. Mössbauer during his graduate studies at Heidelberg in 1957. Mössbauer effect

has provided a significant contribution to the solution of a variety of problems in many

branches of science, especially in the field of solid state chemistry and solid state physics

[198-203].

The hyperfine interactions between a nuclear and appropriate electronic property

are of great significance giving information about the electronic and spin structure. There

are three main hyperfine interactions determined from the Mössbauer spectra are:

(i) The shift in nuclear energy levels induced by the charge of static local

atomic electrons (center shift).

(ii) The energy level splitting in a nucleus due to magnetic fields induced by the

atomic electrons in the vicinity (nuclear Zeeman effect).

(iii) The shift in the levels of nuclear energy induced by a strong electric field

gradient due to the quasi electron (quadrupole splitting).

These characteristics are effects caused by interactions of the absorbing nucleus

with its surroundings. Before the discovery of the Mössbauer effect, the hyperfine

interaction parameters could be inferred indirectly. But with the discovery of the

Mössbauer effect it became

directly with a lot of new information.

Figure 2.5: The Chemical Shift (

levels and the corresponding

The 57Co source is used to probe

decays to 57

Fe by emitting a

most common element examined with Mössbauer spectroscopy

sufficiently abundant,

which is the requirements for the

capture to an excited state of

excitation 1 = 3 23 → 1as shown in the diagram given above (

A Mössbauer spectrometer (

toward and away from the sample by a Mössbauer drive, a collimator to filter the

sample, detector, amplifier, analyzer and a data acquisition system.

Mössbauer absorption spectroscopy, the source is accelerated through a range of

velocities using a Mössbauer drive

became possible to study the hyperfine interactions

of new information.

The Chemical Shift (CS) and Quadrupole Splitting (QS

corresponding Mössbauer spectrum.

Co source is used to probe 57Fe in iron containing samples because

Fe by emitting a γ-ray of the right energy to be absorbed by

most common element examined with Mössbauer spectroscopy because its

has a low energy γ-rays, and a long lifetime

the requirements for the Mössbauer spectrum observed. The

capture to an excited state of 57Fe, after an intermediate step, is subjected to

3 1 1 23 (ground state), which results in 14.4

the diagram given above (Fig. 2.5).

A Mössbauer spectrometer (Fig. 2.6) system consists of a γ-

toward and away from the sample by a Mössbauer drive, a collimator to filter the

sample, detector, amplifier, analyzer and a data acquisition system.


Mössbauer drive to produce a Doppler effect and scan

58

possible to study the hyperfine interactions specifically and

QS) of nuclear energy

Fe in iron containing samples because 57Co

ray of the right energy to be absorbed by 57

Fe. Fe is the

because its isotope 57

Fe is

lifetime excited nuclear state

The 57

Co decays by beta

, is subjected to the de-

14.4 keV Mössbauer line

-rays source oscillating

toward and away from the sample by a Mössbauer drive, a collimator to filter the γ-rays,

sample, detector, amplifier, analyzer and a data acquisition system. During the


to produce a Doppler effect and scan γ-ray energy

59

through a given range which might be ±11 mm/s (1 mm/s = 48.075 neV) [204]. Its most

common form is Mössbauer absorption spectroscopy in which a solid sample is exposed

to a beam of γ-rays and a detector measure the intensity of the beam transmitted through

the sample. During the scanning of γ-rays by Doppler shifting, the detector records the

energy of γ-rays absorbed by the sample. The output of the Mössbauer experiment is the

plot of energy of γ-rays being transmitted through the sample and the source velocity in

mm/s.

Figure 2.6: A schematic view of typical Mössbauer spectrometer [205].

In the present study, Mössbauer analysis was carried out using SEECo MSCI

Mössbauer spectrometer running in constant acceleration mode with a source of 50 mCi

57Co in Rh matrix. The model of Lorentzian multiplet analysis was used to analyze the

data obtained. This model is the standard one permitting to set several Lorentzian

singlets, doublets or sextets, corresponding to paramagnetic sites with or without a

quadrupole splitting and sites with a magnetic hyperfine field and the center shift within

the first order perturbation limit.

Analyzer Data Acquisition

System

Amplifier

60

2.3.4 Magnetic properties measurement systems

Magnetic measurements have been performed by using superconducting quantum

interference device magnetometer and vibrating sample magnetometer at an applied

external field of 50 kOe and 5 kOe, respectively.

2.3.4.1 SQUID magnetometer

Superconducting quantum interference device (SQUID) magnetometer is the most

advanced set to study the magnetic properties of small test samples in a wide range of

temperatures and magnetic fields. It is extremely sensitive to all kinds of AC- and DC-

magnetic measurements. Extremely sensitive magnetic measurements are made with

superconducting pickup coils and a Superconducting Quantum Interference Device

(SQUID). For this reason, this class of instruments is called a SQUID magnetometer

[206-208].

The SQUID magnetometer is a very sensitive instrument for magnetic

measurements that can measure the magnetic fields of the order of femtotesla (10-15 T)

and magnetic moments of the order of 10-10

Am2 [209]. The automated control and data

storage are done by a computer and other electronic controllers. The components of a

SQUID magnetometer are shown in Fig. 2.7.

Figure 2.7: Basic scheme of MPMS Quantum Design [209].

The SQUID magnetometer mainly composed of five parts described below:

(i) Superconducting magnet (solenoid made of superconducting wire).

(ii) The detection coil that detects changes in the external magnetic field and

transforms it into an electric current.

61

(iii) An input coil that converts the resulting current into a magnetic flux in the

SQUID sensor.

(iv) The electronic control which converts the applied flux into a room

temperature voltage output.

(v) The acquisition of hardware and software to acquire, store and analyze data.

The entire structure is composed of liquid Helium Dewar with temperature 4.2 K as

shown in Fig. 2.8.

Figure 2.8: SQUID detection diagram [210].

Data can be collected between H = O to ± 50 kOe and T = 1.7-400 K. The

maximum sensitivity of the instrument is of the order of 10-9

emu. Mass of samples used

is usually 20 to 40 mg but strong magnetic materials can be characterized with less

quantity. The most common tasks to be undertaken for measurements are summarized in

the following form:

i) Mounting sample

ii) Loading sample

iii) Centring sample: The sample must be centered in the SQUID pickup coils to

make sure that all four coils can sense the magnetic moment of the sample. If

the sample is not centered, the coils might be capable to sense the partial

magnetic moment of samples.

62

iv) Measuring sample: When performing an immediate measurement mode,

MPMS MultiVu performs the measurements of the sample under the present

conditions of the system without waiting for any condition to stabilize. If you

want to ensure that conditions are stable when the system starts the

measurement, run the measurement in a sequence and use commands of

appropriate sequence to stabilize the system conditions.

In the present study symmetric magnetic hysteresis loops were measured at 300,

200 and 100 K using a superconducting quantum interference device (SQUID)

magnetometer over a field range of -50 kOe to +50 kOe. The magnetic loops measured at

above mentioned temperatures were used to obtain the basic magnetic parameters namely

saturation magnetization, MS, coercive field, HC, and the magnetocrystalline anisotropy

constant, K1, for all the samples. For calculation of saturation magnetization and

magnetocrystalline anisotropy constant, the high field regions of the M-H loops were

modeled using the Law of Approach to saturation (LoA) [211] while HC was directly

determined by fitting the data near M = 0 kA/m in the M-H loops. The Law of Approach

is said to be valid in the magnetized region 0.97MS<M≤MS [212] therefore, the

corresponding high field region of the magnetic loops was used to determine the

magnetocrystalline anisotropy coefficient, K1. According to the LoA, it was assumed that

as the magnetization approached saturation, all irreversible hysteretic processes were

completed and that the magnetization process in that region was due to reversible rotation

of magnetization against anisotropy and forced magnetization, which can be described by

[211]:

6 67 91 − ;< −

-< −⋯> + ?@ (2.7)

Where ‘a’ and ‘b’ are the fitting coefficients, MS and H are the saturation magnetization

and applied field, respectively, κH is the forced magnetization term. The coefficient ‘a’

relates to domain wall pinning and therefore in the high field region of reversible rotation

of magnetization, the value of a ≈ 0. For randomly oriented polycrystalline samples with

a cubic crystal structure, the coefficient ‘b’ is given as [213]:

A = $BCD E

µF%G (2.8)

63

Where µo is the permeability of free space. The constant 8/105 is specific to cubic

anisotropy of randomly oriented polycrystalline materials. These considerations give the

law of approach used here for analyses:

6 = 67 H1 − $BCD I E

µF%G<JK + ?@ (2.9)

Data from the region above 97% of maximum magnetization value were fitted to

the LoA (Eq. 2.9) to determine the values of parameters MS, K1, and κ.

Figure 2.9: A symmetric hysteresis loop of magnesium ferrite measured by the SQUID

magnetometer.

2.3.4.2 Vibrating sample magnetometer (VSM)

VSM is a powerful technique to study the magnetic properties of materials easily,

reliably and accurately, as the movement of the sample allows discrimination of the

background signals. VSM is widely used to measure the magnetic properties for a wide

range of sample sizes and configurations, i.e., powders, solids, single crystals, thin films,

and liquids, and also it is quite suitable for measurements at low, high, and ambient

temperatures. In the development of nano-ferrite materials the VSM has been widely

used to measure the saturation, remanence, coercivity, anisotropy fields, etc., and also to

-200

-150

-100

-50

0

50

100

150

200

-50 -40 -30 -20 -10 0 10 20 30 40 50M (

kA

/m)

H (kOe)

64

measure the temperature dependent parameters of interest, for example, Curie

temperature.

In a vibrating sample magnetometer (VSM), the sample is subjected to undergo

vibration that induces an electrical signal into a set of suitably placed pick-up or detection

coils (Fig. 2.10). The VSM is composed of a head assembly for the vibration of the drive,

the quartz sample probe, electromagnet assembly, controller and a computer [214].

Figure 2.10: Flow sheet for the working of vibrating sample magnetometer (VSM) [215].

Before the measurements, VSM has to be calibrated with a standard Ni to give a

saturation magnetization of 55 ± 0.5 Am2/kg at 0.5 T. For this, Ni standard is placed at

the geometric center inside the pickup coils to perform a procedure known as "saddling"

as shown in Fig. 2.11. In a standard set of electromagnets, the saddle point is located by

finding the local extrema of the moment signal in the X, Y, and Z planes. In general, the

magnetic moment is set to be a maximum in X and Z directions and minimum along the

Y direction during saddling. The samples for analysis are placed with the same geometry

like that of nickel standard.

65

Figure 2.11: Moment signals to find out the saddle point of the sample [214].

The scan rate of temperature also affects the Curie temperature measurements.

High sweep rate of temperature may yield invalid results. Figure 2.12-A shows the results

of a half-hour measurement of emu versus temperature for a Nickel sample. It can be

seen that the rapid temperature sweep causes the curve to flatten and fall almost vertically

from about 400 °C. This is not valid measurement. Figure 2.12-B shows much better

results by repeating the same measurement at a slow temperature sweep rate of 2-hour.

Hence, for good results, temperature sweep rates should be slow. The steepest change in

magnetic moment causes significant accuracy in measurements.

Figure 2.12: Demagnetization curves of nickel sample as a function of temperature at

two temperature sweep speeds [214].

66

In this study, author measured the temperature dependence of normalized moment

of small chunks of the samples in an applied field of 5 kOe within a temperature range of

350 K to 973 K, using a vibrating sample magnetometer (VSM, Model 7400 Lake Shore,

USA). The sweep rate of temperature for the measurements carried out in the aforesaid

temperature range was almost 8 hrs for the data to be more reliable. The obtained thermo-

magnetic curves were used to find out the Curie temperature, TC, of the samples.

2.3.5 Electrical and dielectric measuring systems

Electrical properties have been studied in terms of measurements of DC-electrical

resistivity, and dielectric properties. The electrical resistivity was measured in a

temperature range 298-673K by using two probes method, while the dielectric properties

measurements were performed at room temperature using an LCR meter bridge (Wayne

Kerr LCR 4275) in a frequency range of 100 Hz to 3 MHz.

2.3.5.1 DC-electrical resistivity

Ferrites have a higher resistance than the metals by several orders of magnitude

and are considered highly sensitive to the structure of materials. The two-probe method

was used to measure the resistivity of Mg-ferrite nanomaterials in the temperature range

300-675K. As the resistivity of the Mg-ferrites is very high and it was not feasible to use

the four probe method, because it is suitable for materials with low resistance [36] and

also not suitable for high temperature measurements due to instability of silver paste used

for connections. Two probes method was developed for measuring electrical properties

such as resistivity, mobility and the activation energy at high temperature (300-675 K).

The main components of this device are given below:

1) High voltage power (up to 210 V)

2) A heating element

3) A thermocouple

4) A control TRAIC

5) Digital multimeter

6) A sample holder equipped with two electrodes

67

Above mentioned components are assembled as described below (Fig. 2.13).

Sample holder equipped with two electrodes which are arranged symmetrically with

respect to each other on opposite sides of the sample, was inserted into a small heating

oven that was fitted inside a hollow ceramic cylindrical jar. The shape of the sample

analyzed in two probe method can be disc like, cylindrical, cube, or parallelepiped. A

thermocouple connected with a UT-55 multimeter was capable of measuring the

temperature in the range 73 to 1473 K. A triac controller was inserted in series with the

heater to heat the sample at a controlled rate. DC-power supply and SourceMeter

(KEITHLEY 2400) was connected in series with the sample holder. The temperature was

noted with a step of 5 K, and the corresponding change of current flowing through a

sample is measured by the SourceMeter (KEITHLEY 2400) with a given voltage applied

across the sample.

Heater

V

A

Sample

220 V

Traic Control

Figure 2.13: Flow sheet diagram of the two-probe resistivity apparatus [216].

In the present work DC-electrical resistivity was measured with disk shape pellets

of the sample having 13mm diameter and 2mm thickness. The electrical resistance of a

material is defined as the ratio between the applied voltage (V) to electrical current (I)

flowing through it, given by Ohm's law:

L = /1 (2.10)

The resistivity of the samples has been calculated by following equation [217]:

N = LO/P (2.11)

68

Where, R is resistance, L is the thickness and A is the cross sectional area of the pellets

(of the dimensions mentioned on the previous page). The variation of electrical resistivity

with temperature measured by the above mentioned device is shown in Figure 2.14.

Figure 2.14: Temperature dependence of electrical resistivity measured by the device

described in Sec. 2.3.5.1.

The activation energy and drift mobility of all the samples were also calculated

from DC-electrical resistivity data. The electrical resistivity of the ferrite materials

decreases with the increase in temperature obeying the Arrhenius type equation [34]:

N = NQRST 9 UVWX> (2.12)

Where, kB is the Boltzmann constant, T is temperature (K) and Ea is the activation energy

needed for the hopping of an electron from one ion to the other neighboring ion of the

same element with the different valence state, which give rise to a conduction in

materials.

The drift mobility (µd) of the doped Mg-ferrite samples was calculated using the

relation [34]:

YZ = B[!\ (2.13)

0

1

2

3

4

5

6

7

8

275 325 375 425 475 525 575 625 675

ρ (

Ω.c

m)1

09

T(K)

69

Where, e is the charge of the electron, ρ the DC-electrical resistivity at given temperature

and n is the concentration of charge carriers which can be calculated from the relation:

= '(Z]^_+% (2.14)

Where, NA is the Avogadro’s number, db the bulk density, M the molecular mass of the

sample and pFe is the number of iron atom in the chemical formula of the samples.

Figure 2.15: Arrhenius type relationship for activation energy calculation

2.3.5.2 Dielectric properties measuring system

Spinel ferrites are attractive materials with many applications in high frequency

microwave devices. Therefore, it is important to study their dielectric behavior at

different frequencies. The inductance, capacitance and resistance (LCR) Meter Bridge

provide economical, thorough and accurate testing with a basic accuracy of 0.1%, of any

dielectric materials at frequencies up to several frequency ranges. It can automatically

measure the inductance (L), capacitance (C) and resistance (R) and other parameters such

as the quality factor, impedance, AC resistance, conductance and dielectric loss tangent

of the subjected materials.

It consists of a power supply (PS) to generate the AC voltage, resistor, and

amplifiers. Voltage is applied to the sample (S) through the source resistance 'RS' (B),

which varies depending on the measuring range. The current flows through the resistor

12

14

16

18

20

22

1.4 1.6 1.8 2.0 2.2 2.4 2.6

ln ρ

103/T

70

RR to operational amplifiers (A1 and A2). The amplifier automatically adjusts the

voltage gain level, so that an electric current passing through a resistor should be equal to

the current flowing through the sample and provides an output signal proportional to that

current. The voltage across the sample is measured by a separate signal path with a four-

wire Kelvin connection. Out of these four wires Kelvin, two wires used to perform the

test current and two independent wires for the detection of voltage across the sample. The

real and imaginary signals are obtained by multiplying the voltage and current signals

with a reference signal in phase with supply voltage and other shifted 900 from the power

supply voltage. These signals are read by a microprocessor (M). This prevents the voltage

drop in the current carrying wires to affect the voltage measurement. The LCR meters are

controlled by a high speed microcontroller that runs the display, keypad, general purpose

interface bus (GPIB), computer interfaces and handler interface. The schematic diagram

of LCR meter is shown in Fig. 2.16.

Figure 2.16: Schematic diagram of an LCR meter [218].

The dielectric constant (έ) and dielectric loss tangent (tanδ) are calculated from

data of inductance, capacitance and resistance. The Dielectric constant is a measure of the

degree to which a medium can resist the flow of charge, is calculated as follows:

έ = `aℰcd (2.15)

Where, C is the capacitance of the pellet in farad, L the thickness and A the cross-

sectional area of the pellet and ℰo the constant of permittivity of free space. The amount

of energy loss during each cycle is expressed as the dielectric loss tangent (tanδ) of

Microprocessor

(M)

/Sample

(S)

(B)

(PS)

71

dielectric material and its value is directly given by the instrument in the form of D. D is

called the dissipation factor and is the real part of the impedance divided by the reactance

(the imaginary part of the impedance). The change in dielectric constant (έ) with applied

field frequency measured by the above mentioned apparatus is shown in Fig. 2.17.

Figure 2.17: Frequency dependence of the dielectric constant measured by the apparatus

described in Sec. 2.3.5.2.

0

4

8

12

16

20

24

4 6 8 10 12 14 16

έ(1

03)

ln f

CHAPTER 3

RESULTS AND DISCUSSION

72

3 RESULTS AND DISCUSSION

3.1 Thermal Properties

Thermal analysis (TGA/DTG) of the as-synthesized magnesium ferrite and its

metal doped derivatives has been performed to evaluate the mechanism of phase

formation of a spinel cubic lattice and to assess the thermal events on heating the samples

up to 1250 K.

Figure 3.1 shows the TGA/DTG curves for magnesium ferrite (MgFe2O4)

precursor. During heat treatment of the sample, dehydration, decomposition and removal

of any interfering phase take place. TGA curve depicts multiple weight loss regions.

Initial weight loss (~ 7 %) between room temperature to 450 K, as indicated by first

minima in the corresponding DTG curve, is attributed to evaporation of the residual

water. The second sharp minima at 548 K, corresponds to the weight loss (~ 10 %) due to

decomposition of metal hydroxides into their corresponding oxides. In this temperature

range, the NH4NO3 may also dissociate to liberate NOx and O2 by the following reaction

[188]:

NH4NO3 → NOx + O2 + H2O (3.1)

However, a minor weight loss (obvious from DTG curve) observed up to the

temperature of 800 K is due to complete decomposition of any interfering phases, or

attributed to the formation of an intermediate hematite (α-Fe2O3) and mono-ferrite

(MFe2O4) phases [177]. Above this temperature, negligible weight loss indicated that the

thermal events ceased beyond this temperature. Although the desired cubic spinel

structure formation has been started earlier at 900 K but complete single-phase is formed

at higher temperature of 1123 K at which the intermediate phases are completely

removed, as can be seen from XRD analysis in the next Sec. 3.2.1. Hence, an annealing

temperature of 1123 K has been selected for stable metal oxide phase formation of all the

samples.

73

Figure 3.1: TGA/DTG curves of the as-synthesized pure magnesium ferrite

The TGA/DTG curves of the samples doped with binary mixture of Co-Cr, Ni-Cr,

Cu-Cr, Zn-Cr and Mn-Cr are shown in Figs. 3.2-3.6. One representative sample with a

composition of x = 0.4 from each of the five series is selected for thermal analysis. It can

be seen from Figs. 3.2-3.6, that the addition of the substitutes has a slight influence on the

thermal behavior of the doped magnesium ferrites. The Co-Cr, Zn-Cr and Mn-Cr series

show larger weight loss than that of the undoped magnesium ferrite. This weight loss is

due to evaporation of the bivalent dopants i.e. Co, Zn and Mn, at high temperature [219].

However, TGA curves of the doped samples are similar to that of the undoped sample

(Fig. 3.1). All samples exhibit dehydration process up to 400 K, followed by

decomposition of metal hydroxides and nitrates in the range of 500-600 K. The

decomposition of metal hydroxides (M(OH)2 where, M = Co, Ni, Cu, Zn and Mn) usually

takes place in two steps [220]:

M(OH)2 → MOOH + H2O (3.2)

MOOH→ M2O3 +H2O (3.3)

-0.015

-0.010

-0.005

0.000

0.005

0.010

0.015

80

83

86

89

92

95

98

101

350 500 650 800 950 1100 1250

mg (

%)

K-1

Weig

ht (%

)

T (K)

TGA

DTG

74

The metal hydroxides of the doped samples decomposed into corresponding metal

oxides by heat treatment up to 600 K, followed by the removal of any interfering phase.

Then rearrangement of these metal oxides into mixed oxides occurred through a phase

transformation at temperature above 900 K, which has no significant weight loss

indicating the beginning of the formation of the cubic spinel phase. However, the

formation of pure single phase of magnesium ferrite and its derivatives required high

annealing temperature for the removal of any interfering phases. Thermal behavior of the

samples also shows that the conversion of precursor into crystalline metal oxide is a

kinetically controlled process by means of a thermal decomposition of the metal

hydroxides.

Figure 3.2: TGA/DTG curves of the as-synthesized Mg0.6Co0.4Cr0.4Fe1.6O4

-0.006

-0.004

-0.002

0.000

0.002

0.004

0.006

65

70

75

80

85

90

95

100

320 470 620 770 920 1070 1220

mg (

%)

K-1

Weig

ht (%

)

T (K)

TGA

DTG

75

Figure 3.3: TGA/DTG curves of the as-synthesized Mg0.6Ni0.4Cr0.4Fe1.6O4

Figure 3.4: TGA/DTG curves of the as-synthesized Mg0.6Cu0.4Cr0.4Fe1.6O4

-0.010

-0.006

-0.002

0.002

0.006

0.010

80

84

88

92

96

100

330 480 630 780 930 1080 1230

mg (

%)

K-1

Weig

ht (%

)

T (K)

TGA

DTG

-0.010

-0.005

0.000

0.005

0.010

0.015

0.020

75

80

85

90

95

100

300 450 600 750 900 1050 1200

mg (

%)

K-1

Weig

ht (%

)

T (K)

TGA

DTG

76

Figure 3.5: TGA/DTG curves of the as-synthesized Mg0.6Zn0.4Cr0.4Fe1.6O4

Figure 3.6: TGA/DTG curves of the as-synthesized Mg0.6Mn0.4Cr0.4Fe1.6O4

-0.004

-0.003

-0.002

-0.001

0.000

0.001

0.002

0.003

0.004

60

65

70

75

80

85

90

95

100

300 450 600 750 900 1050 1200

mg (

%)

K-1

Weig

ht (%

)

T (K)

TGA

DTG

-0.003

-0.002

-0.001

0.000

0.001

0.002

0.003

65

70

75

80

85

90

95

100

300 450 600 750 900 1050 1200

mg (

%)

K-1

Weig

ht (%

)

T (K)

TGA

DTG

77

3.2 Structural Properties

Structural analyses of the synthesized magnesium ferrites and its substituted

derivatives have been performed by X-ray diffraction (XRD), scanning electron

microscopy (SEM) and energy dispersive X-ray fluorescence (ED-XRF) to evaluate the

crystal structure, surface morphology and elemental composition of the materials,

respectively.

3.2.1 X-ray diffraction (XRD) analysis

Well-defined diffraction peaks i.e. (220), (311), (400), (422), (511) and (440) are

observed in the sample which corresponds to magnesium ferrite (Fig. 3.7). The observed

peaks closely correspond to the standard pattern for cubic spinel magnesium ferrite

(ICDD-01-073-1720). The results of XRD analysis are in conformity with the TG

analysis (Fig. 3.1) described before in Sec. 3.1.

Figure 3.7: XRD patterns of the synthesized MgFe2O4 sample compared with the

standard pattern.

2θ (degree)

Inte

nsity (

%)

0

50

100

20 30 40 50 60 70 80

Ref. Pattern: Magnesium diiron (iii) oxide, 01-073-1720

220

31

1

400

422

511

44

0

MgFe2O4

78

XRD patterns of the Co-Cr doped Mg1-xCoxCrxFe2-xO4 samples (x = 0.0-0.5) are

illustrated in Fig. 3.8. The diffraction peaks closely correspond to the standard pattern

(Fig. 3.7, ICSD Ref. No. 01-073-1720) of a spinel magnesium ferrite with CCP structure

and shows no unidentified peaks which imply that the Co-Cr, in the selected substitution

level, can completely be accommodated in the cubic spinel lattice and no extra phase

such as hematite (Fe2O3) is present. The crystallite size (D), calculated by the well-known

Debye-Scherrer formula [214], is found to be in the range 23-47 nm (Table 3.1). A

significant decrease of ~24 nm in particle size is noted in response to the increase in Co-

Cr substitution levels. Moreover, the calculated crystallite size for Co-Cr substituted

magnesium ferrites is smaller as compared to that reported (40–75 nm) for the spinel

ferrite [222] synthesized by the conventional ceramic route. It has been reported that a

crystallite size of < 50 nm is desirable for obtaining a suitable signal-to-noise ratio for

switching applications [223].

Lattice constant (a), and cell volume (Vcell) are calculated from XRD data by

using Eqs. 2.3-2.4. Data of the Table 3.1 show that there is a minor decrease in the value

of lattice constant and cell volume which might be due to smaller ionic size of the doped

Cr3+

(0.63 Å) ions compared to that of the host ion i.e. Fe3+

(0.64 Å). While, the doped

Co2+ (0.72 Å) ion has an ionic size equal to that of host Mg2+ (0.72 Å) ion and therefore,

have no influence on the lattice constant.

X-ray density (dx) and bulk density (db) are calculated using the Eqs. 2.5-2.6. The

value of dx increases from 4.50 g/cm3

to 4.87g/cm3

with the increase in Co-Cr contents

due to the larger molar mass of the doped metal cations as shown in Table 3.1. Data of

the Table 3.1 show that the magnitude of bulk density, db for the synthesized samples in

present study, is smaller than that of dx (db is ~88 % of dx) as anticipated due to some

unavoidable pores created during sintering.

79

Figure 3.8: XRD patterns of Mg-ferrite doped with Cox-Crx contents (x = 0.0-0.5).

2θ (degree)

Inte

nsity (

a.u

)

0.0

0.1

0.2

0.3

0.4

0.5

30 40 50 60 70

22

0

311

40

0

51

1

42

2 44

0

80

Table 3.1: Crystallite size (D), lattice constant (a), cell volume (Vcell), X-ray density (dx)

and bulk density (db) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5)

Parameters x =0.0 x=0.1 x=0.2 x=0.3 x=0.4 x=0.5

D (nm) 47 40 30 35 23 28

a (Å) 8.384 8.380 8.380 8.379 8.378 8.379

Vcell (Å3) 589 588 588 588 588 588

dx (g/cm3) 4.50 4.59 4.66 4.73 4.80 4.87

db (g/cm3) 3.93 4.04 4.11 4.17 4.24 4.31

The XRD patterns for the series doped with Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr are

shown in Figs. 3.9-3.12. The XRD peaks for the synthesized samples and corresponding

Miller indices (hkl) exactly match with the standard pattern of cubic spinel magnesium

ferrite [ICDD-01-073-1720] assuring the single phase formation of the synthesized

materials. This confirms that all the dopants must have simply replaced the Mg2+

-Fe3+

ions without distortion of the cubic symmetry of the host magnesium ferrite except for

higher levels of Cu-Cr substitution which causes slight tetragonal distortion in the cubic

symmetry of the lattice. The crystallite sizes lie in the ranges of 15-31, 39-62, 25-36 and

27-40 nm for the Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr substituted magnesium ferrites,

respectively (Tables 3.2-3.5). All the crystallite sizes are small enough to obtain the

suitable signal to noise ratio for switching applications.

The value of lattice constant (a) and cell volume (Vcell) decreased with the

increase in Ni-Cr, Cu-Cr and Mn-Cr co-substitution level while both parameters

increased for the Zn-Cr substituted series. This variation in both parameters is more likely

a compensatory response of the crystal structure to accommodate the dopants of different

ionic size as compared to those of host ions i.e. Mg2+

(0.72 Å) and Fe3+

(0.64 Å). As we

can see that in case of Ni-Cr and Mn-Cr series, ionic size of the divalent dopants i.e. Ni2+

(0.69 Å) and Mn2+ (0.67 Å) is smaller as compared to that of host Mg2+ (0.72 Å) ions

which causes the shrinkage of crystal lattice that leads to decrease in the values of the

lattice constant and cell volume, while Cr3+

ions having ionic size (0.63 Å) comparable to

81

Fe3+

(0.64 Å) which is being replaced, would show negligible influence on these

parameters. On the other hand, an increase in the above-mentioned parameters with the

increase in Zn-Cr substitution level, is owed to the influence of the ionic size of doped

metal cations. The ionic size of divalent dopant ions i.e. Zn2+

(0.82 Å) is a greater than

that of replacing Mg2+

(0.72 Å) ion and caused an increase in lattice constant and cell

volume. In case of Cu-Cr substituted series, the dopants ions (Cu2+

(73 Å), Cr3+ (0.63 Å))

have comparable ionic size with those of host ions i.e. Mg2+

(0.72 Å) and Fe3+

(0.64 Å).

On the contrary, the lattice constant and cell volume both do not have the amenable

decreasing trend (Table 3.3) with the increase in the Cu-Cr concentration. Thus, the most

likely cause of such behavior could be the tetragonal distortion in the cubic lattice

structure.

The observed increase in the value of X-ray density (dx) and bulk density (db) with

the incorporation of dopant ions for all synthesized series (Tables 3.1-3.5) is due to (i)

Larger molar mass of the doped metal cations compare to that of host pair i.e. Mg-Fe and

(ii) shrinkage of lattice that leads to the observed decrease in the value of Vcell.

82

Figure 3.9: XRD patterns of Mg-ferrite doped with Nix-Crx contents (x = 0.0-0.5).

2θ (degree)

0.0

0.1

0.2

0.4

0.5

0.3

30

Inte

nsity (

a.u

)

40 50 60 70

22

0 3

11

40

0

42

2

511

440

83

Figure 3.10: XRD patterns of Mg-ferrite doped with Cux-Crx contents (x = 0.0-0.5).

0.5

0.3

0.1

0.0

30 40 50 60 70

2θ (Degree)

Inte

nsity (

a.u

)

22

0 31

1

400

422

511

44

0

0.4

0.2

84

Figure 3.11: XRD patterns of Mg-ferrite doped with Znx-Crx contents (x = 0.0-0.5).

2θ (Degree)

Inte

nsity (

a.u

)

0.0

0.1

0.2

0.3

0.4

0.5

220

31

1

40

0

44

0

511

422

30 40 50 60 70

85

Figure 3.12: XRD patterns of Mg-ferrite doped with Mnx-Crx contents (x = 0.0-0.5).

2θ (Degree)

Inte

nsity (

a.u

)

0.0

0.1

0.2

0.3

0.4

0.5

22

0

31

1

40

0

440

51

1

42

2

30 40 50 60 70

86


and bulk density (db) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5).


D (nm) 47 31 26 18 18 15

a (Å) 8.384 8.383 8.382 8.380 8.378 8.375

Vcell (Å3) 589 589 589 588 588 587

dx (g/cm3) 4.50 4.58 4.65 4.72 4.80 4.87

db (g/cm3) 3.93 4.02 4.10 4.16 4.21 4.24


and bulk density (db) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5).


D (nm) 47 39 44 54 56 62

a (Å) 8.384 8.365 8.360 8.357 8.357 8.349

Vcell (Å3) 589 585 584 584 584 582

dx (g/cm3) 4.50 4.64 4.71 4.79 4.86 4.95

db (g/cm3) 3.93 4.11 4.19 4.26 4.35 4.43

87


and bulk density (db) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5).


D (nm) 47 25 32 36 25 27

a (Å) 8.384 8.389 8.395 8.399 8.406 8.412

Vcell (Å3) 589 590 592 592 594 595

dx (g/cm3) 4.50 4.58 4.66 4.73 4.81 4.88

db (g/cm3) 3.93 4.01 4.08 4.16 4.26 4.33


and bulk density (db) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5).


D (nm) 47 27 28 40 28 35

a (Å) 8.384 8.380 8.377 8.375 8.372 8.370

Vcell (Å3) 589 588 588 587 587 586

dx (g/cm3) 4.50 4.58 4.64 4.70 4.77 4.83

db (g/cm3) 3.83 4.03 4.08 4.14 4.22 4.28

3.2.2 Scanning electron

The scanning electron micrograph of pure

in Fig. 3.13. SEM micrograph clearly shows that the

be an agglomeration of the individual particles. The shaded areas in the micrograph are

due to the voids. The presence of voids on the sample surface shows their porous nature,

as has also been concluded on the basis of comparison of the X

(Table 3.1-3.5). This indicates

the synthesized samples leading to the low value of bulk density

exhibits well-defined crystalline

coalescence to form bigger

Figure 3.13: Scanning electron micrograph of

The surface micrographs

and Mn-Cr are shown

and 0.3 are taken from each

general, it is evident that

spherical with a uniform surface morphology.

exhibits well-defined crystalline nanoparticles.

of smaller particles and appear

lectron microscopic (SEM) analysis

The scanning electron micrograph of pure magnesium ferrite (

micrograph clearly shows that the surface is less smooth and appears to


The presence of voids on the sample surface shows their porous nature,

as has also been concluded on the basis of comparison of the X-

). This indicates that the annealing process induced unavoidable po

amples leading to the low value of bulk density

defined crystalline nanoparticles of spherical shape that resulted due to their

bigger particles during annealing at high temperat

Scanning electron micrograph of MgFe2O4 sample

icrographs of the samples doped with Co-Cr

are shown in Figs. 3.14-3.18. Samples with dopant composition

en from each of the five series due to limited facility

general, it is evident that the majority of the particles of the synthesized samples

spherical with a uniform surface morphology. The surface of the synthesized samples

fined crystalline nanoparticles. Few of them are formed

and appeared as lumps. The grain growth of such agglomerates on the

88

ferrite (MgFe2O4) is shown

surface is less smooth and appears to


The presence of voids on the sample surface shows their porous nature,

-ray and bulk densities

that the annealing process induced unavoidable pores in

amples leading to the low value of bulk density. The sample surface

that resulted due to their

particles during annealing at high temperature.

Cr, Ni-Cr, Cu-Cr, Zn-Cr

compositions of x = 0.1

series due to limited facility of SEM analysis. In

the synthesized samples is

The surface of the synthesized samples

formed as agglomerates

of such agglomerates on the

89

surface is however, normal. The individual nanoparticles may tend to coalesce together

due to annealing of the samples at high temperature of 1123 K during their synthesis.

The particle size in each of the synthesized series, gets smaller when the dopant

content increases except for Cu-Cr doped series (Figs. 3.14-3.18), as also evident from

the XRD data (Tables 3.1-3.5). Figures 3.13-3.18 reflect that the particle size appears to

be larger for the Cu-Cr doped samples [177] as compared to the rest of the doped

samples. Further, the reduction in particle size of the doped samples (Figs. 3.14-3.18) as

compared to the undoped sample (Figs. 3.13) is owing to the incorporation of the co-

dopants in the magnesium ferrite nanomaterials.

Since the average particle size of all the doped samples (except Cu-Cr co-doped

magnesium ferrites) is below 50 nm and it can be useful to obtain a suitable signal-to-

noise ratio for switching applications [34], as described previously.

Figure 3.14: Scanning electron micrographs of Mg1-xCoxCrxFe2-xO4 (x = 0.1, 0.3).

x = 0.1 x = 0.3

90

Figure 3.15: Scanning electron micrographs of Mg1-xNixCrxFe2-xO4 (x = 0.1, 0.3).

Figure 3.16: Scanning electron micrographs of Mg1-xCuxCrxFe2-xO4 (x = 0.1, 0.3).

x = 0.3 x = 0.1

x = 0.1 x = 0.3

91

Figure 3.17: Scanning electron micrographs of Mg1-xZnxCrxFe2-xO4 (x = 0.1, 0.3).

Figure 3.18: Scanning electron micrographs of Mg1-xMnxCrxFe2-xO4 (x = 0.1, 0.3).

In conclusions, the spinel cubic phase formation begins at ~800 K and completes

at 1123 K as it is evident from the TGA/DTG and XRD analyses. Doping of a binary

mixture of metals retains the cubic geometry of the crystal lattice, with an average

crystallite size of 15-47 nm. There is the strong influence of the substituents on different

parameters, i.e. lattice constant, cell volume, X-ray density and bulk density, of

nanocrystalline materials. The substitution of cationic mixtures in magnesium ferrite has

x = 0.1 x = 0.3

x = 0.3 x = 0.1

92

resulted in the homogenous surface morphology with the mixtures of individual

nanoparticles and their aggregates. Thus, the results obtained from XRD and SEM

analyses indicate the formation of nanosized crystallites in the doped Mg-ferrites.

3.2.3 Energy dispersive X-ray fluorescence (ED-XRF) analysis

Elemental compositions of some representative samples of the each synthesized

series with composition x = 0.2 and 0.4, revealed that the elemental compositions of the

samples are in conformity to their nominal stoichiometry (Table 3.6). Only cation

compositions have been shown in the table because ED-XRF is insensitive to oxygen.

Therefore, it has been assumed that the oxygen concentration is the same as that in the

chemical formulae.

Table 3.6: The observed and theoretical composition of selective samples of

Mg1-xMxCrxFe2-xO4 (M = Co, Ni, Cu, Zn, Mn, and x = 0.0, 0.2 and 0.4).

Theoretical values Observed values (molL-1

)

Mg M Cr Fe

MgFe2O4 0.97 ----- ----- 2.02

Mg0.8Co0.2Cr0.2Fe1.8O4 0.76 0.20 0.19 1.83

Mg0.6Co0.4Cr0.4Fe1.6O4 0.58 0.38 0.38 1.61

Mg0.8Ni0.2Cr0.2Fe1.8O4 0.79 0.19 0.20 1.83

0.58 0.43 0.39 1.62 Mg0.6Ni0.4Cr0.4Fe1.6O4

Mg0.8Cu0.2Cr0.2Fe1.8O4 0.75 0.21 0.19 1.81

0.57 0.42 0.41 1.63 Mg0.6Cu0.4Cr0.4Fe1.6O4

Mg0.8Zn0.2Cr0.2Fe1.8O4 0.77 0.20 0.18 1.81

0.56 0.39 0.40 1.63 Mg0.6Zn0.4Cr0.4Fe1.6O4

Mg0.8Mn0.2Cr0.2Fe1.8O4 0.78 0.22 0.18 1.83

0.57 0.44 0.38 1.62 Mg0.6Mn0.4Cr0.4Fe1.6O4

93

3.3 Mössbauer Analysis

Mössbauer spectra of the Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) samples recorded at

room temperature (300 K) are illustrated in Fig. 3.19. The magnetic components of the

spectra have been fitted by using Lorentzian line fitting to assess the number of

interactions with a minimum χ2~1.0 and solid lines in Fig. 3.19 show the obtained

computer fitted Lorentzian curves. All samples exhibit well resolved and magnetically

normal Zeeman split sextets attributed to the presence of iron ions at both tetrahedral (A-

sites) and octahedral (B-sites) sites which indicates the ferrimagnetic nature of the

synthesized materials with mixed spinel structure [224]. However, a slight broadness in

Zeeman lines is observed as evident from Fig. 3.19. The broadening of lines could be

attributed to slight changes in the magnetic environment surrounding Fe3+

ions in the

same sub lattice; such changes in environment of Fe3+

ions would lead to a change in the

magnetic field and consequently results in some appreciable broadening of the Zeeman

lines [225]. In addition, the changes in the magnetic environment would affect center

shift (CS) values slightly, displacing the sextets with respect to one another and cause

general broadening of these lines. The presence of relaxation effect also conceivably

results in broadening of the lines. The increased broadening in the lines might be due to

larger probable distribution of Co2+

and Cr3+

ions surrounding Fe3+

ions at B-sites.

A summary of the various hyperfine interaction parameters is presented in Table

3.7 and obtained results are in close agreement with the studies reported earlier for spinel

ferrites [226, 227]. The center shift (CS) results from the electrostatic interaction between

the charge distribution of the nucleus and S-electrons with finite probability being found

in the region of the nucleus. Data of the Table 3.7 show that the center shift (CS (A) and

CS (B)) values have random trend, indicating that S-electron charge distribution of Fe3+

is

not much influenced by Co-Cr substitution. As expected, the outer sextet shows a larger

center shift (∆δ~0.04-0.10 mm/s) because of the difference in the Fe3+

-O2-

inter-nuclear

separation normally larger for B-site ions as compared to that for A-site ions.

Consequently, smaller overlapping between orbitals of Fe3+

and O2-

at B-site occurs,

resulting in smaller covalency and hence the larger center shift, CS, for B-site Fe3+

ions

[226]. The sextets belonging to octahedral (B) and tetrahedral (A) sites have been

assigned on the basis of the values of hyperfine magnetic field (H) at the nucleus and

94

center shift (CS). In general, most of the ferrites have higher values of hyperfine magnetic

field and center shift correspond to B-site sextet and the lower values of the same

correspond to A-site sextet [198].

Figure 3.19: Mössbauer spectra of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5)

-20 10 -10 0 20

Velocity (mm/s)

Rela

tive I

nte

nsity

0.5

0.4

0.3

0.2

0.1

0.0

95

The values of quadrupole splitting (refer Table 3.7) for hyperfine spectra of all the

samples have found to be negligibly small and can be attributed to the fact that overall

cubic symmetry is not much altered between Fe3+ ions and their surroundings by

substitution with Co-Cr ions in magnesium ferrites.

Table 3.7: Center shift (CS), quadrupole splitting (QS), hyperfine magnetic field (H) and

relative area (A) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5).

Parameters Bond Area x=0.0 x=0.1 x=0.2 x=0.3 x=0.4 x=0.5

CS (mm/s) A-site 0.286 0.263 0.309 0.318 0.299 0.300

B-site 0.383 0.370 0.344 0.361 0.354 0.345

QS (mm/s) A-site -0.002 0.008 0.005 -0.046 0.007 0.015

B-site -0.121 -0.120 -0.141 -0.031 -0.114 -0.017

H (kOe) A-site 536.3 527.8 514.3 511.8 504.9 498.4

B-site 558.2 549.3 546.7 537.2 519.1 518.3

A (%) A-site 30.6 40.1 47.0 51.5 55.1 64.2

B-site 69.4 59.9 53.0 48.5 44.9 35.8

The variation in hyperfine magnetic field, (HA & HB), at the two sublattices, A-

and B-sites with increase of Co-Cr contents is shown in Table 3.7. The hyperfine

interaction at A-site is found to decrease from 536.3 to 498.4 kOe, whereas at B-site it

decreases from 558.2 to 518.3 kOe with the increase in Co-Cr co-substitution level. The

magnetic environment of iron ions at both the sites has been changed by substitution with

Co-Cr ions, thus accounting for the decrease in the value of HA and HB. The larger

hyperfine field assigned to the B-sites is attributed to the dipolar field result of the

covalent nature of the tetrahedral bonds [228]. The center shift and a larger hyperfine

field for the same site helped in assigning the outer sextet to the octahedral sites and the

inner sextet to the tetrahedral sites. The A-site hyperfine fields are typically 4-5% less

than that of B-site and the difference is usually attributed to the larger covalency of Fe3+

-

96

O2-

at the A-site. The variation in hyperfine fields is due to the change in A-B and B-B

super-exchange hyperfine interactions, as the cation neighbors around Fe3+

ions are

changed due to substitution of Co-Cr contents into Mg-ferrite [79].

The relative area (A) under the resonance curve of the sub-spectra deduced from

the measurements is helpful to conceive the Co, Cr and Fe site occupancy in sub-lattice

sites. The relative peak-area of B-site decreases with the incorporation of Co-Cr and

subsequently observed to increase for that of A-site. The incorporation of Co-Cr shows

variation in the peak-area of Mössbauer sub-spectra which determines the fractions of

iron ions at A- and B-sites. This variation in relative peak-area is supported by

Mössbauer spectra illustrated in Fig. 3.19, that the intensity of outer sextet decreases

while that of inner sextet increases with the incorporation of Co-Cr ions. Thus, we can

conclude that Co2+

ions occupy the octahedral sites with Cr3+

ions [37, 229] and the

system possesses a partial inverse spinel structure.

Mössbauer spectra of the Mg1-xMxCrxFe2-xO4 (M = Ni, Cu, Zn, Mn and x = 0.0-

0.5) samples recorded at room temperature (300 K) are illustrated in Figs. 3.20-3.23. The

magnetic components of the spectra have been fitted by using Lorentzian line fitting to

assess the number of interactions with a minimum χ2~1.0. All the samples exhibit

magnetically normal Zeeman split sextets which indicates the ferrimagnetic nature of the

samples except for the Zn-Cr substituted samples with compositions x = 0.4 and 0.5

which show less spectra splitting with broad sextets. One small doublet also appears at x

= 0.5 which might be due to the magnetically isolated Fe3+

ions located at one of the two

spinel sites which could not participate in the long-range magnetic ordering due to

reduction in Fe3+ ions on either site and large number of nonmagnetic nearest neighbors.

For diamagnetically substituted ferrites, the existence of a central doublet superimposed

on less-resolved magnetic sextets has been reported for a number of systems [230, 231].

A summary of the hyperfine interaction parameters for Ni-Cr, Cu-Cr, Zn-Cr and

Mn-Cr substituted series is presented in Tables 3.8-3.11. These results are in good

agreement with the literature for spinel ferrites [226, 227]. Data of the Tables 3.8-3.11

show that the center shift CS (A) and CS (B) values are fluctuating randomly independent

of the dopant concentration indicating that S-electron charge distribution of Fe3+

is not

much influenced by co-substitution. As expected, the center shift for the B-site Fe3+

ions

97

is larger than that for the A-site Fe3+

ions because of relatively larger Fe3+

–O2-

bond

separation in the former one.

It is also obvious from Tables 3.8-3.11 that a very small quadrupole splitting (QS)

for A- and B-sites have been observed in all the samples. This suggests that the

coexistence of negligible chemical disorder and overall cubic symmetry causes no net

quadrupole shifts in Zeeman sextets [225].

The variation in hyperfine magnetic fields for A- (HA) and B-sites (HB) with M-Cr

(Ni, Cu, Zn and Mn) concentration (x = 0.0-0.5) is shown in Tables 3.8-3.11. The

hyperfine interaction at A- and B-site is found to decrease as the substitution level of co-

dopant increases from x = 0.0 to x = 0.5. The A-site hyperfine fields are typically 4-16%

less than that of B-site and the difference is usually attributed to the larger covalency of

Fe3+

–O2-

at the A site as discussed before. The variation in hyperfine fields is due to the

change in A-B and B-B super-exchange hyperfine interactions as the cation neighbors

around Fe3+

ions are changed [79].

By concluding from the above discussion it becomes evident that the values of

center shift (CS) and quadrupole splitting (QS) show that the S-electron density is not

much disturbed by the substitution of these dopants. The hyperfine interaction (H) and

relative area (A) at the two sites show that all the substituents occupy the octahedral site

except Zn2+

which shows a strong preference for tetrahedral site.

98

Figure 3.20: Mössbauer spectra of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5)

-20 10 -10 0 20

Velocity (mm/s)

Rela

tive I

nte

nsity

x = 0.5

x = 0.4

x = 0.3

x = 0.2

x = 0.1

x = 0.0

99

Fig. 3.21: Mössbauer spectra of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5).

-20 20 0 10 -10

Velocity (mm/s)

Rela

tive I

nte

nsity

0.5

0.4

0.3

0.2

0.1

0.0

100

Fig. 3.22: Mössbauer spectra of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5).

Rela

tive I

nte

nsity

Velocity (mm/s)

-20 -10 0 10 20

0.5

0.4

0.3

0.2

0.1

0.0

101

Fig. 3.23: Mössbauer spectra of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5).

Rela

tive I

nte

nsity

Velocity (mm/s)

-20 -10 0 10 20

0.5

0.4

0.3

0.2

0.1

0.0

102


relative area (A) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5).

Parameters Lattice site x=0.0 x=0.1 x=0.2 x=0.3 x=0.4 x=0.5

CS (mm/s) A-site 0.286 0.263 0.266 0.296 0.299 0.300

B-site 0.383 0.360 0.381 0.392 0.377 0.384

QS (mm/s) A-site -0.002 0.068 0.017 0.031 0.007 0.010

B-site -0.121 -0.130 -0.124 -0.118 -0.114 -0.112

H (kOe) A-site 536.3 528.2 511.6 504.7 504.9 507.8

B-site 558.2 547.5 533.7 518.6 521.1 524.5

A (%) A-site 30.6 40.1 45.0 57.2 56.4 50.5

B-site 69.4 59.9 55.0 42.8 43.6 48.5


relative area (A) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5).


CS (mm/s) A-site 0.286 0.261 0.273 0.279 0.274 0.281

B-site 0.383 0.371 0.367 0.363 0.372 0.379

QS (mm/s) A-site -0.002 0.013 -0.017 0.003 -0.001 -0.111

B-site -0.121 -0.106 -0.002 -0.109 -0.110 0.006

H (kOe) A-site 536.3 516.2 489.5 468.2 441.7 423.4

B-site 558.2 524.4 509.1 493.6 490.3 488.7

A (%) A-site 30.6 35.3 38.3 42.9 47.5 51.3

B-site 69.4 64.7 61.7 57.1 52.5 48.7

103

Table 3.10: Center shift (CS), quadrupole splitting (QS), hyperfine magnetic field (H)

and relative area (A) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5).


CS (mm/s) A-site 0.286 0.335 0.308 0.317 0.326 0.329

B-site 0.383 0.366 0.336 0.318 0.337 0.362

QS (mm/s) A-site -0.002 0.018 -0.010 -0.032 0.003 0.016

B-site -0.121 0.008 0.006 0.019 0.003 0.012

H (kOe) A-site 536.3 485.5 458.1 437.4 427.9 405.7

B-site 558.2 544.2 533.1 507.3 471.1 484.4

A (%) A-site 30.6 39.8 44.9 49.4 49.7 51.5

B-site 69.4 60.2 54.1 50.6 50.3 49.5

Table 3.11: Center shift (CS), quadrupole splitting (QS), hyperfine magnetic field (H)

and relative area (A) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5).


CS (mm/s) A-site 0.286 0.305 0.306 0.297 0.286 0.249

B-site 0.383 0.356 0.342 0.331 0.339 0.329

QS (mm/s) A-site -0.002 0.027 -0.023 -0.023 0.003 0.023

B-site -0.121 -0.002 -0.006 -0.009 -0.007 -0.001

H (kOe) A-site 536.3 525.5 503.9 497.4 483.0 471.1

B-site 558.2 544.6 529.2 527.3 524.2 523.7

A (%) A-site 30.6 38.3 49.8 52.3 47.8 61.6

B-site 69.4 61.7 50.2 47.7 42.2 38.4

104

3.4 Magnetic Measurements

3.4.1 SQUID magnetometric measurements

The magnetic hysteresis loops are measured by superconducting quantum

interference device (SQUID) magnetometer at three different temperatures of 300, 200

and 100 K to find out the magnetic parameters such as saturation magnetization (MS),

remanence (Mr), magnetocrystalline anisotropy coefficient (K1) and coercivity (HC). The

effects of temperature and composition on these parameters have been investigated.

Figure 3.24 shows the symmetric magnetic hysteresis loops of Co-Cr substituted

magnesium ferrite nanomaterials of composition; Mg0.8Co0.2Cr0.2Fe1.8O4 at 300, 200 and

100 K. The inset shows the first quadrant of magnetic hysteresis loops measured at 300

K. Saturation magnetization (MS), remanence (Mr), coercivity (HC), and first-order cubic

anisotropy coefficient (K1) for all samples are extracted from the hysteresis loops.

For calculation of saturation magnetization and magnetocrystalline anisotropy

coefficient, the high field regions of the M-H loops (Fig. 3.24) are modeled using the Law

of Approach (LoA) to saturation [211] based on the assumption that at sufficiently high

field only the rotational processes remain with an additional forced magnetization term

that is linear with the applied field. The values of Mr and HC have been determined

directly by fitting the data at H = 0 kOe and M = 0 kA/m in the M-H loops, respectively.

Magnetocrystalline anisotropy is an important performance property in evaluating

magnetic materials for different applications. In order to evaluate the temperature and

composition dependence of anisotropy, it is assumed that all irreversible hysteretic

processes have been completed when the major hysteresis loop closed and that a further

increase of the magnetic moment is due to the rotational processes which are connected

with the magnetic anisotropy. Based on the Law of Approach (LoA) to saturation, which

describes the dependence of magnetization (M) on the applied field for H»HC,

magnetization near the saturation magnetization (MS) can be experimentally obtained as

described in Sec. 2.3.4.1 (Eq. 2.9).

105

Figure 3.24: M-H loops of Mg0.8Co0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset: First

quadrant of magnetic hysteresis loops for Mg1-xCoxCrxFe2-xO4 at 300 K.

Data from the region above 97% of maximum magnetization value have been

fitted to the LoA (Eq. 2.9) and the values of MS and K1 at 300, 200 and 100 K are

determined. However, in the case of measurements carried out at 100 K, detailed

examination of the magnetic loops revealed that these measured loops could not approach

to complete saturation even at high field of µoH = 50 kOe. So, MS and K1 were the only

fitting parameters as magnetization is negligible in this regime at this temperature, as

might be expected, and, therefore, κ = 0 has been taken at 100 K.

Data of the Table 3.12 show that the saturation magnetization and remanence

increased monotonically with decreasing temperature from 300 to 100 K for all Co-Cr

doped samples. The values of MS, computed by fitting Eq. 2.9 to the experimental data,

are found to be approximately the same as that measured maximum magnetization values

at µoH = 50 kOe for 300 and 200 K. However at 100 K, the measured value of maximum

magnetization is larger than the MS value by several percent due to the applied field no

longer being able to saturate these samples at the lower temperature. On addition of Co-

106

Cr contents, the MS value initially increases from 148, 171 and 192 kA/m (x = 0.0) to

299, 371 and 415 kA/m (x = 0.3) at 300, 200 and 100 K, respectively, but then decreases

for higher level of substitution (Table 3.12). The Mr value is also varied in a manner

similar to that of MS as discussed above. It increases from 26, 37 and 49 kA/m (x = 0.0)

to 85, 255 and 344 kA/m (x = 0.3) at 300, 200 and 100 K, respectively and afterward

tends to decrease for higher level of substitution (Table 3.12).

Table 3.12: Saturation magnetization (MS), remanence (Mr), magnetocrystalline

anisotropy coefficient (K1) and coercivity (HC) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) at

300, 200 and 100 K.


Ms (kA/m)

@300 K

148 169 218 299 258 187

Mr (kA/m) 26 51 84 85 85 58

K1 (J/m3)10

5 2.31 2.77 3.79 4.04 3.02 2.51

HC (Oe) 57 275 344 265 262 162

Ms (kA/m)

@200 K

171 200 253 371 332 238

Mr (kA/m) 37 113 180 255 220 162

K1 (J/m3)10

5 2.36 4.66 6.33 6.23 5.96 4.56

HC (Oe) 84 970 1384 1186 1099 965

Ms (kA/m)

@100 K

192 228 291 415 376 281

Mr (kA/m) 49 157 226 344 308 225

K1 (J/m3)10

5 2.45 6.27 8.38 7.19 7.12 6.44

HC (Oe) 100 2455 3685 3491 3153 3101

This observed variation in the values of MS and Mr can be understood by

considering the structure of the ferrimagnetic cubic spinels. The number of the octahedral

sites (B-sites) are twice the number of tetrahedral sites (A-sites) and magnetic moments

on the octahedral and tetrahedral sites are antiparallel to each other, thus leading to a net

magnetic moment, M=Moct-Mtet. Since, Mg2+ is a non-magnetic ion; it has no contribution

107

in the magnetic moment of the mixed spinel ferrite (MgFe2O4) which is thus entirely due

to the uncompensated spins of the un-evenly distributed iron ions at both A- and B-sites.

In Co-Cr co-substituted a magnesium ferrite system, Mg and Fe ions are partially

distributed over both sites (A- and B-site) and the doped Co-Cr ions are believed to

substitute preferably into the B-sites [232-234]. Replacement of non-magnetic Mg2+

ion

by magnetic ion (Co2+

) would thus enhance the magnetic moment of B-site leading to an

increase in the net magnetic moment. On the other hand, substitution of magnetic Fe3+

ion (5µB) by less magnetic Cr3+ (3µB) ion would result in dilution of magnetization at B-

site. This competition leads to an increase in the value of MS and Mr to a certain level of

substitution and to a decrease afterwards. However, the overall effect of Co-Cr

substitution is the increase in MS value which could be suitable for different magnetic

applications [224].

Similarly, the temperature and composition dependence of the calculated

anisotropy coefficient (K1) for different Co-Cr contents is shown in Table 3.12. K1 is

observed to increase in magnitude monotonically with decrease in temperature. The

analysis of the temperature dependence of cubic anisotropy of Co-Cr substituted

magnesium ferrite can be divided into two temperature zones, (i) relatively high

temperature (300 K and 200 K) and (ii) low temperature (100 K). For the first zone, the

maximum applied field is sufficiently large compared to the anisotropy field, which can

be estimated as = 2/ [235]. This allows successful application of LoA. As the

temperature decreases, the ratio of exchange interaction to thermal energy increases, that

contributes to the increase in anisotropy for the Co-Cr doped samples studied here. Thus,

at low temperature (the second zone in our case), the anisotropy is so high that it prevents

a complete approach to saturation even at the highest applied field of µoH = 50 kOe (see

Fig. 3.24). At such temperature, the use of the LoA is questionable and, therefore, the

anisotropy coefficient K1 might be suspected even if calculated with the force

magnetization coefficient set to zero, i.e., κ = 0 [236]. The results of Shenker [237], who

determined the cubic anisotropy of CoFe2O4 using single crystals and torque

measurements near their easy axes, support our assumption of high anisotropy fields

discussed above.

108

The calculated values of anisotropy coefficient (K1) for different Co-Cr contents

show that K1 increases substantially to the order of magnitude of 105 J/m

3 for Co-Cr

composition, x ≤ 0.3 but decreases for higher level of Co-Cr contents. The composition

dependence of anisotropy of Mg1-xCoxCrxFe2-xO4 can be interpreted in terms of the

effects of the substituents on site occupancies of the cations. The obtained results suggest

that the substitution of Co+2

for magnesium and Cr3+

for iron; both substituents prefer to

be incorporated into the octahedral sites. According to the one-ion model, the presence of

Co2+ ions on the octahedral sites of the spinel structure enhances the anisotropy of ferrites

[238]. This point of view supports the observed increase in anisotropy coefficient with

increasing Co2+

content up to a certain level of substitution. However, in case of higher

level of substitution, the effect of Cr3+

substitution becomes more prominent and leads to

a reduction of magneto-crystalline anisotropy, as Cr3+

has a strong preference for the

octahedral sites and may even displace some of Co2+ ions to tetrahedral sites [233, 234].

By increasing the Co-Cr substitution level into MgFe2O4, the HC value increases

initially to the dopant content, x < 0.2, and then it continued to decrease for higher level

of substitution as shown in Table 3.12. The value of HC also increases with the decrease

in temperature. The increase in HC value is more prominent at lower temperature i.e. 100

K. It is well known that as the measurement temperature fall down beyond the Curie

temperature of magnetic materials, the materials become more and more anisotropic and

cause an enhancement in the coercive field of the subjected materials.

Figures 3.25-3.28 show the symmetric magnetic hysteresis loops of one

composition of each synthesized series of substituted magnesium ferrite nanomaterials as

an example, i.e. Mg0.8M0.2Cr0.2Fe1.8O4 (M = Ni, Cu, Zn and Mn) at 300, 200 and 100 K.

First quadrant of magnetic hysteresis loops of Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr

substituted magnesium ferrite nanomaterials measured at 300 K are shown in the inset of

their respective Figs. 3.25-3.28. The measured M-H loops were modeled using the Law of

Approach (LoA) to saturation as discussed before.

Data of the Tables 3.13-3.16 show that the values of saturation magnetization and

remanence both increased monotonically by decreasing the temperature from 300 to 100

K. The values of MS, computed by fitting Eq. (2.9) to the experimental data, are found to

be approximately the same as that measured by using maximum magnetization values at

109

50 kOe for measurement temperature of 300, 200 and 100 K. Data of the Tables 3.13-

3.16 show the composition dependence of saturation magnetization and remanence for

Ni-Cr, Cu-Cr and Zn-Cr substituted series. The values of MS and Mr decrease up to a

substitution level of x = 0.3 but above this specific level, both values tend to increase for

Ni-Cr and Zn-Cr series, while continued to decrease for Cu-Cr substituted series.

However, the values of MS and Mr for Mn-Cr doped samples (refer Tables 3.12 and 3.16)

varied in a manner similar to that of Co-Cr doped samples with the incorporation of the

dopants in the magnesium ferrite. The variation in the observed values of MS and Mr can

be interpreted in terms of the site occupancies of the cations in the cubic spinel lattice of

substituted magnesium ferrite.

The observation that with low substitution levels of Ni-Cr (x = 0.0-0.3), the

decreases in the value of MS indicates that initially Ni-Cr ions are substituted into the B-

site which results in a decrease in the net magnetic moment compared to pure magnesium

ferrite. At higher substitution levels, the net magnetization increases significantly because

Ni is being increasingly substituted into the A-sites. Also, as the number of Mg2+

ions at

B-sites decreases, some of the Fe3+

ions could be displaced from A- to B-sites (supported

by Mössbauer data) to compensate the reduction of Mg2+

ions. Hence, the higher

substitution level of Ni-Cr (x = 0.4 and 0.5) would result in the enhancement of the

magnetic moment of the B-site than that of A-site, the net magnetic moment (M = ΣMB-

ΣMA) thus increases.

The variation in the observed values of MS and Mr with Cu-Cr substitution (Table

3.14) can also be explained on the basis of exchange interaction between ions at

tetrahedral (A) and octahedral (B) sites. As it has been mentioned before, the Mg2+ and

Fe3+

ions occupy both A- and B-sites, while the Cu2+

[177] and Cr3+

ions have a

preference for the B site in the Mg-Cu-Cr ferrite. Replacement by magnetic Cu2+

ion

(1µB) for non-magnetic Mg2+

ion (0µB) would help to enhance the magnetization of B-

site. On the other hand, Cr3+

ions (2.81µB) replace Fe3+

ions (5 µB) on the octahedral site

causing a decrease of the magnetic moment of this sub-lattice resulting in a decrease of

the net magnetic moment and consequently, magnetization. Hence, the overall effect of

Cu-Cr co-substitution is the gradual reduction in MS and Mr.

110

The variation in the observed values of MS and Mr for Mg1-xZnxCrxFe2-xO4 (Table

3.15) can be explained in terms of the cation site occupancies. It would be expected that

the more the cation distribution (tends towards mixed spinel), the higher the measured net

moment. Consequently, at 100 K and 200 K, samples with composition x = 0.0 has lower

net magnetization (Ms) than that of the sample with dopant content x = 0.1. On the other

hand, at 300 K the net MS value is enhanced due to the reason that cation site occupancy

is temperature dependent. For Zn-Cr substitution level i.e. x =0.2 and 0.3, Ms value

decreases with increase in dopant content at all temperatures. This might be due to the

replacement of non-magnetic Mg2+

ion (0 µB) at tetrahedral sites by nonmagnetic Zn2+

ion

(0 µB) [239] that has no effect on the magnetic moment of the A-sites. On the other hand,

substitution of Cr3+

ions (2.81 µB) for Fe3+

ions (5 µB) on the octahedral site reduces the

total magnetic moment of the sub-lattice. At x > 0.3, the increase in Ms could be due to

the increase in Zn2+ substitution for Mg2+at the A-sites, which resulted in some Fe3+ being

displaced from the A-sites to the B-sites. As Zn2+

ions are incorporated at the expense of

Mg2+

ion, higher level of Zn2+

substitution would continue to occupy A-site [240] while

the number of Mg2+

ions decreased at B-site. Consequently, some Fe3+

shifted from A-

site to B-site to compensate the reduction in the number of Mg2+

ions at B-site. Since the

net magnetization (MS) is given by the difference between ΣMB-sites and ΣMA-sites, the

displacement of Fe3+

from A- to B-site would increase ΣMB-sites (but decreases ΣMA-sites),

hence MS would increase. The observation that both Ms and Mr follow a similar trend may

suggest that a similar mechanism is responsible for their variation.

As can be seen from Table 3.16, that on addition of Mn-Cr contents into

magnesium ferrite, the saturation magnetization, MS, and remanence, Mr, appear to

increase up to certain doping level (for x ≤ 0.3) and then decreases for higher contents of

dopants (x > 0.3) at all the temperatures. This variation in the value of MS and Mr can be

understood by conceiving that the Mg and Fe ions are partially distributed over both A

and B sites while Mn-Cr ions substitute predominantly into the B-sites in the magnesium

ferrite [233, 241]. Replacement of non-magnetic Mg2+ ion by magnetic ion (Mn2+) in the

magnesium ferrite system would enhance the magnetic moment of B-site leading to an

increase in the net magnetic moment. On the other hand, substitution of Fe3+

(5µB) by less

magnetic Cr3+

(3µB) ion would result in the dilution of magnetization. Hence, owing to a

111

competition between the two ions to occupy the same site, leads to an increase in the

value of MS and Mr up to a certain level of substitution and to decrease afterwards.

The temperature and composition dependence of the first-order cubic anisotropy

coefficient (K1) for different series of Mg1-xMxCrxFe2-xO4 (M = Ni, Cu, Zn and Mn) is

shown in Tables 3.13-3.16. The temperature and composition dependence of the

calculated anisotropy coefficient (K1) for Ni-Cr contents are shown in Table 3.13. The

value of K1 increases in magnitude with the decrease in temperature. This is because

when the measurement temperature decreases away from the Curie temperature of the

samples, the ratio of the exchange interaction to thermal energy increases, which

contributes towards the increase in anisotropy. The variation in the value of the

anisotropy coefficient (K1) with Ni-Cr substitution is in the range of 1.25×105-3.69×10

5

J/m3

as shown in Table 3.13. The value of K1 varies in a way similar to that of MS value

except for x = 0.5, which decreases rather than increase like MS value. A close

observation of the high regions of M-H loops shows that, for a specimen with x ≤ 0.4, the

applied field is high enough to overcome anisotropy, hence causing a complete approach

to saturation. The first-order cubic anisotropy coefficient (K1) changed with composition

similar to MS value. The deviation observed at the substitution levels above x = 0.4, can

be due to anisotropy field higher than the maximum applied field of the µoH = 50 kOe

that prevented a complete approach to saturation. For this, the forced magnetization

constant is set to zero, so that the calculations are made by the Eq. 2.9, with MS and K1

being the only fitting parameters. However, the assumption of complete approach to

saturation of the LoA method is not fulfilled in this case, the calculation of K1, although

indicative, is not considered accurate and somewhat smaller than the expected value [34].

The temperature and composition dependence of the calculated anisotropy

coefficient (K1) for Cu-Cr contents are shown in Table 3.14. The temperature dependence

of cubic anisotropy of Cu-Cr substituted magnesium ferrite can be interpreted by

considering that the maximum applied field is large enough compared to the anisotropy

field, which can be estimated as = 2/ [235], and this allows us to apply the

LoA successfully. K1 has been observed to increase in magnitude with the decrease in

temperature as explained before.

112

The composition dependence of the calculated anisotropy coefficient (K1) for

different Cu-Cr contents shows that the anisotropy coefficient is of the order of

magnitude 105 J/m3, it increased only for initial doping level (x = 0.1) and afterwards

continued to decrease with increasing Cu-Cr contents. This is thought to be due to a

reduction of the tetrahedral-octahedral exchange coupling. The substitution of Cu-Cr

reduces the exchange coupling between the octahedral and tetrahedral sites. Probably this

reduction in exchange coupling is responsible for the reduction in the magnitude of

magnetic anisotropy, despite the fact that upon Cu-Cr substitution, the amount of the

magnetic Cu2+

ions (1µB) in the octahedral sites most likely replaces the non-magnetic

Mg2+

ions (0µB). On the other hand, the Cr3+

ions (2.81µB) replace the highly magnetic

Fe3+

ions (5µB) which overtakes the magnetic effect of Cu2+

(1µB) and results in the

weakening of exchange interactions [177]. In case of Cu-Cr substituted compositions of

Mg1-xCuxCrxFe2-xO4, the complete saturation is brought about by lower anisotropy and

makes the validity of the Law of Approach to saturation at all measurement temperatures

i.e. 300, 200 and 100 K.

The temperature and composition dependence of the first-order cubic anisotropy

coefficient (K1) for different Zn-Cr contents are shown in Table 3.15. In this case K1 is

also observed to increase with the decrease in temperature like that of Co-Cr, Ni-Cr and

Cu-Cr series discussed before. For Zn-Cr substituted series i.e. Mg1-xZnxCrxFe2-xO4, the

complete saturation is brought about by lower anisotropy and makes the validity of the

Law of Approach (LoA) to saturation at lower temperature of 100 K.

The composition dependence of K1 has similar behavior to that of MS, suggesting

that cation distribution is responsible for the observed variation. At 200 and 100 K, the

value of K1 increases with Zn-Cr substitution level up to x = 0.1 which indicates that the

exchange interaction initially increases due to doping of Zn-Cr contents. On the contrary,

a similar trend is not observed at 300 K and might be due to the effect of increased

temperature on cation distribution. However, K1 decreases on further substitution of Zn-

Cr from x = 0.1 to x = 0.3 because the increase in Zn-Cr concentration reduces the

exchange interaction between the octahedral and tetrahedral sites. The reason behind the

observed increase in K1 for x > 0.3 might be that the further substitution of Zn-Cr resulted

in stronger exchange coupling due to the displacement of some Fe3+

from A-site to B-

113

site. The substitution of Zn-Cr altered the A-B exchange coupling as can be seen from the

steep decrease in Curie temperature with the dopant contents up to x = 0.3 in the next sec.

3.4.2. It is probable that this reduction in exchange coupling is responsible for the

observed decrease in magnetic anisotropy coefficient.

The temperature and composition dependence of the calculated anisotropy

coefficient, K1 for Mn-Cr contents are shown in Table 3.16. The value of K1 is observed

to vary in a manner similar to that of Co-Cr substituted series. This can be explained on

the bases of the same mechanism as discussed before for Co-Cr substituted series. It

increases in magnitude monotonically with decrease in temperature. While, the

composition dependence of the calculated anisotropy coefficient (K1) for different Mn-Cr

contents shows that K1 increases substantially with increasing Mn-Cr contents that are of

the order of magnitude 105 J/m

3, for the specimen with x ≤ 0.3 and decreases for higher

level of Mn-Cr contents.

It has been found that in certain cases of low temperature, the anisotropy of

samples is so high that it prevents a complete approach to saturation even at maximum

applied field of 50 kOe [34, 242]. However, the anisotropy results are consistent with the

reported experimental work and theoretical predictions for spinel ferrites [237, 243].

The temperature and composition dependence of coercivity (HC) for different

compositions of Mg1-xMxCrxFe2-xO4 (M = Ni, Cu, Zn and Mn) is shown in Tables 3.13-

3.16. By increasing the dopant level, the HC value increases initially up to the dopant

content of x ≤ 0.1, and then it decreases for higher dopant concentration of Ni-Cr and Cu-

Cr doped samples as shown in Tables 3.13 and 3.14, respectively. The value of HC also

increases with the decrease in temperature. This increase in HC is more prominent at

lower temperature, 100 K. It is well known that as the measurement temperature go

beyond the Curie temperature of magnetic materials, the materials become more and

more anisotropic [244] that cause an enhancement in coercivity of the subjected

materials. In case of Ni-Cr doped samples, the value of coercivity at 300 K is found to be

in the range of 43-127 Oe. Both anisotropy field and small grain boundary volume

introduced during sample preparation can influence coercivity. Anisotropy does not seem

to be the dominating factor in deciding the coercivity of the Ni-Cr substituted series.

However, in regards to the influence of grain size on coercivity, it is known that there is

114

no domain wall in the single-domain grains and the magnetization mechanism is a

domain rotation process [34]. As a result, the HC is in direct proportion to the volume of

single-domain grains; it decreases with as the single-domain particle size decreases.

It is well known that coercivity of polycrystalline ferrites is strongly dependent

upon the magnetocrystalline anisotropy and grain size. In Cu-Cr substituted magnesium

ferrites, the value of first-order anisotropy coefficient is small and has no dominating

influence on the coercivity. In polycrystalline materials, regarding correlation between

coercivity and grain size, it is expected that larger grains will provide less pinning of

domain walls because of the lower volume fraction of grain boundaries. Accordingly, we

conjecture that the pinning of magnetization at the grain boundaries is the most likely

cause for the variation in coercivity. Basically, the results obtained for Cu-Cr substituted

magnesium ferrites are consistent with previous reports, from which it is evident that the

coercivity of ferrites depends on the temperature and the grain size [245, 246], indirectly

on chemical substitution of polycrystalline materials. Table 3.3 shows that the sample

with x = 0.1 contains grains of smaller size compared to those of the sample with the

dopant content of x = 0.5, since the coercivity has the opposite trend with grain size: the

coercivity is reduced by an increase in Cu-Cr contents. Therefore, we can conclude that

the coercivity is reduced due to increasing Cu-Cr level in the prepared samples.

The coercivity shows a gradual decrease with Zn-Cr substitution up to x ≤ 0.3,

except for some initial increase from x = 0.1 to x = 0.2 at 100 K. Afterwards, a slight

increase is observed for x > 0.3. Coercivity also increases with decrease in temperature as

shown in Table 3.15. The observed trend in the variation of coercivity might be related to

the variation in magneto-crystalline anisotropy and small grain boundary volume mainly

introduced by synthetic routes.

The values of coercivity (HC) as a function of Mn-Cr content is shown in Table

3.16. It is clear from the Table 3.16 that the coercivity decreases continuously from 57 to

25 Oe with Mn-Cr substitution level. Its value increased progressively with the decrease

in temperature (refer Table 3.16) from 300 to 100 K. This decreasing trend with Mn-Cr

substitution might be due to decrease in magneto-crystalline anisotropy and small grain

boundary volume. The variation in coercivity is in conformity, to some extent, with the

magneto-crystalline anisotropy.

115

The magnetic recording media requires the saturation magnetization value as high

as possible and moderate coercivity values [247]. In the present study, the saturation

magnetization of magnesium ferrite increases from 148 kAm-1 to 299, 206, 164 and 290

kAm-1

by doping with Co-Cr, Ni-Cr, Zn-Cr and Mn-Cr, respectively while its value

decreases on doping with Cu-Cr from 148 to 55 kAm-1

at most potential operational

range around 300 K. The coercivity is found to decrease with the substitution level for all

the series studied here. The M-Cr (M = Co, Ni, and Mn) substituted samples can be

useful for the applications in the parallel recording media because these samples have the

moderate coercivity in the desired range as stated above and has an enhanced value of the

saturation magnetization.

Figure 3.25: M-H loops of Mg0.8Ni0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset: First

quadrant of magnetic hysteresis loops for Mg1-xNixCrxFe2-xO4 at 300 K.

116

Figure 3.26: M-H loops of Mg0.8Cu0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset: First

quadrant of magnetic hysteresis loops for Mg1-xCuxCrxFe2-xO4 at 300 K.

Figure 3.27: M-H loops of Mg0.8Zn0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset: First

quadrant of magnetic hysteresis loops for Mg1-xZnxCrxFe2-xO4 at 300 K.

117

Figure 3.28: M-H loops of Mg0.8Mn0.2Cr0.2Fe1.8O4 at 300, 200 and 100 K; Inset: First

quadrant of magnetic hysteresis loops for Mg1-xMnxCrxFe2-xO4 at 300 K.

118


anisotropy coefficient (K1) and coercivity (HC) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5) at

300, 200 and 100 K.


Ms (kA/m)

@300 K

148 146 143 129 173 206

Mr (kA/m) 26 22 22 17 28 29

K1 (J/m3)10

5 2.31 3.69 1.58 1.25 3.08 2.26

HC (Oe) 57 127 106 66 51 43

Ms (kA/m)

@200 K

171 183 179 161 216 258

Mr (kA/m) 37 33 32 26 42 44

K1 (J/m3)10

5 2.36 4.43 2.20 1.51 3.70 2.90

HC (Oe) 84 191 159 99 77 65

Ms (kA/m)

@100 K

192 225 220 198 266 317

Mr (kA/m) 49 46 44 36 59 62

K1 (J/m3)10

5 2.45 5.32 3.08 2.10 4.44 3.40

HC (Oe) 100 287 239 149 116 97

119


anisotropy coefficient (K1) and coercivity (HC) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5) at

300, 200 and 100 K.


Ms (kA/m)

@300 K

148 137 116 105 74 55

Mr (kA/m) 26 31 26 18 12 7

K1 (J/m3)10

5 2.31 2.33 2.14 1.25 1.02 0.83

HC (Oe) 57 71 60 46 39 35

Ms (kA/m)

@200 K

171 170 141 139 104 80

Mr (kA/m) 37 51 40 29 22 13

K1 (J/m3)10

5 2.36 2.98 2.37 1.51 1.28 0.95

HC (Oe) 84 121 83 61 54 49

Ms (kA/m)

@100 K

192 212 169 165 129 106

Mr (kA/m) 49 67 54 46 34 22

K1 (J/m3)10

5 2.45 3.50 2.65 1.97 1.47 1.18

HC (Oe) 100 179 128 89 81 77

120


anisotropy coefficient (K1) and coercivity (HC) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) at

300, 200 and 100 K.


Ms (kA/m)

@300 K

148 142 129 118 157 164

Mr (kA/m) 26 30 22 15 27 28

K1 (J/m3)10

5 2.31 2.24 2.15 2.06 2.33 2.51

HC (Oe) 57 55 53 48 55 64

Ms (kA/m)

@200 K

171 177 161 141 207 226

Mr (kA/m) 37 46 41 28 44 48

K1 (J/m3)10

5 2.36 2.57 2.41 2.24 2.59 2.88

HC (Oe) 84 80 78 62 70 73

Ms (kA/m)

@100 K

192 216 206 182 245 288

Mr (kA/m) 49 85 55 39 66 74

K1 (J/m3)10

5 2.45 3.02 2.80 2.31 3.04 3.30

HC (Oe) 100 153 143 89 94 101

121


anisotropy coefficient (K1) and coercivity (HC) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5) at

300, 200 and 100 K.


Ms (kA/m)

@300 K

148 164 269 290 211 149

Mr (kA/m) 26 7 11 15 8 5

K1 (J/m3)10

5 2.31 2.81 3.27 3.61 2.89 2.29

HC (Oe) 57 44 41 40 28 25

Ms (kA/m)

@200 K

171 205 336 363 264 186

Mr (kA/m) 37 11 17 23 12 8

K1 (J/m3)10

5 2.36 3.37 3.92 4.33 3.47 2.75

HC (Oe) 84 66 62 60 42 38

Ms (kA/m)

@100 K

192 252 413 447 325 229

Mr (kA/m) 49 15 24 32 17 11

K1 (J/m3)10

5 2.45 4.04 4.71 5.20 4.16 3.30

HC (Oe) 100 96 90 87 65 50

122

3.4.2 Vibrating sample magnetometric (VSM) measurements

Temperature dependence of the magnetic moment is considered a suitable tool for

the analysis of various kinds of magnetic materials. The magnetic moment is measured as

a function of temperature by using the vibrating sample magnetometer and the Curie

temperature (TC) is determined from the plots between temperature and normalized

moment. In the temperature range below the Curie point, the samples are considered to be

pure ferrimagnetic with all spins arranged collinearly. At temperature exceeding the TC,

the thermal energy of the system becomes too large to disrupt the aligned spins, giving

rise to a complete paramagnetic behavior [248].

The variation of normalized moment, for Co-Cr substituted magnesium ferrite

nanomaterials, with temperature in the range of 350-1123 K at low applied magnetic field

(~5 kOe) is shown in Fig. 3.29. These plots are used to find out the Curie temperature

(TC) of the samples by extrapolating the steep region. In the beginning, magnetic moment

decreases slowly with increasing temperature followed by a sharp decrease on reaching

the Curie temperature (TC) and after this it tends level off.

Figure 3.29: Thermal variation of normalized moment in Mg1-xCoxCrxFe2-xO4

(x = 0.0-0.5).

0.0

0.2

0.4

0.6

0.8

1.0

300 400 500 600 700 800 900 1000

No

rma

lized

mom

en

t

T (K)

0.0 0.1

0.2 0.3

0.4 0.5

123

The Curie temperature determined from the plot between normalized moment and

temperature for MgFe2O4 is 681 K as shown in Table 3.17. The value of TC represents the

characteristic magnetic interactions and overall strength of the exchange interactions

between the spins of the cations (A-B exchange interactions) present at the two sites (A-

and B-sites). The Curie temperature (TC) increases from 681 K for x = 0.0 to maximum

value of 766 K for x = 0.2, and then decreases for x ≥ 0.3 (Table 3.17). This could occur

due to modification of the A-B super-exchange interaction on Co-Cr substitution into

Mg-ferrite. This increase in TC value is due to replacement of non-magnetic Mg2+ ion

(0µB) by a magnetic Co2+

ion (3µB), that would strengthen the A-B super-exchange

interaction, since Co2+

ions are well known to produce large induced anisotropy owing to

their relatively high orbital contribution to the magnetic moment [75]. While, the

decrease in the TC value might be due to higher content (x > 0.2) of less magnetic Cr3+

(3

µB) ions replacing more magnetic Fe3+ (5 µB) ions that in turn could be responsible for the

weakening of the A-B super-exchange interactions.

The variation in a normalized moment of magnesium ferrites nanomaterials doped

with Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr is shown in Figs. 3.30-3.33. The value of Curie

temperature for all the samples determined from these plots is shown in Table 3.17.

For Ni-Cr substituted series TC increases from 681 K to 832 K with the dopant

contents as shown in Table 3.17. An increase in the value of TC could be related to a

number of factors including, the number and the type of magnetic ions present in the two

sub-lattices of the spinel crystal and their exchange interactions. Considering the type of

cation and their exchange interaction, the observed increasing trend of TC does not agree

with the previous observations from Mössbauer data and rest of magnetic parameters

observed for Ni-Cr doped series. However, the increasing trend in the TC value of spinel

ferrites has been reported previously with decrease in the grain size [249, 250] as has

been observed in the present study.

The Curie temperature of the Cu-Cr doped series decreases continuously from 681

K (x = 0.0) to 591 K (x = 0.5) with the increase in dopant content (Table 3.17) and can be

explained on the basis of the number of magnetic ions present in the two sub-lattices and

their mutual interactions. As it is mentioned earlier, the gradual replacement of Fe3+

and

Mg2+

with Cu-Cr substitution causes the reduction in magnetization at B-sites, and results

124

in the weakness of A-B super-exchange interactions. Thus, the thermal energy required to

offset the spin alignment decreases, thereby decreasing the Curie temperature with the

increase of Cu-Cr contents. A similar decrease of TC with the doped impurity ions was

also observed by Rana et al., [69] who investigated Mg-Ni ferrites and by Melagiriyappa

[251] who studied Mg-Zn ferrites.

The Curie temperature of the Zn-Cr doped series decreases from 681 K (x = 0.0)

to 558 K (x = 0.5) with an increase in the dopant substitution levels. This variation in TC

can be qualitatively explained on the basis of assumption; as the Zn-Cr proportion

increases, the relative number of ferric ions on the A- and B-sites diminish that reduces

the strength of A-B super-exchange interactions, thereby decreasing the Curie

temperature of the samples.

Table 3.17: Curie temperature (TC) of Mg1-xMxCrxFe2-xO4 (M = Co, Ni, Cu, Zn, Mn and

x = 0.0-0.5)

Specimen x=0.0 x=0.1 x=0.2 x=0.3 x=0.4 x=0.5

Mg1-xCoxCrxFe2-xO4

TC (K)

681 727 766 682 660 618

Mg1-xNixCrxFe2-xO4 681 728 747 766 786 832

Mg1-xCuxCrxFe2-xO4 681 650 635 619 598 591

Mg1-xZnxCrxFe2-xO4 681 667 641 634 589 558

Mg1-xMnxCrxFe2-xO4 681 697 713 738 649 603

In case of Mn-Cr substituted samples, TC value increased from 681 K (x = 0.0) to

738 K for x = 0.3 and then decreased for x ≥ 0.4 (Table 3.17). This happened due to

substitution of non-magnetic Mg2+

ion (0µB) with magnetic Mn2+

ion (5µB), that would

strengthen the A-B super-exchange interaction owing to increase in TC up to a certain

level of substitution. The decrease in the TC value for x ≥ 0.4, might be due to the higher

content of less magnetic Cr3+

(3 µB) ions replacing more magnetic Fe3+

(5 µB) ions that in

turn could be responsible for the weakening of the A-B super-exchange interactions.

Although, Mn2+ has a similar value of the magnetic moment as that of the Fe3+ yet the co-

doping of Mn-Cr could be responsible for the weakening of interactions.

125

Figure 3.30: Thermal variation of normalized moment in Mg1-xNixCrxFe2-xO4

(x = 0.0-0.5)

Figure 3.31: Thermal variation of normalized moment in Mg1-xCuxCrxFe2-xO4

(x = 0.0-0.5).

0.0

0.2

0.4

0.6

0.8

1.0

300 400 500 600 700 800 900 1000

No

rma

lized

mom

en

t

T (K)

0.0 0.10.2 0.30.4 0.5

0.0

0.2

0.4

0.6

0.8

1.0

300 400 500 600 700 800 900 1000

No

rmaliz

ed

mom

ent

T (K)

0 0.1

0.2 0.3

0.4 0.5

126

Figure 3.32: Thermal variation of normalized moment in Mg1-xZnxCrxFe2-xO4

(x = 0.0-0.5).

Figure 3.33: Thermal variation of normalized moment in Mg1-xMnxCrxFe2-xO4

(x = 0.0-0.5).

0.0

0.2

0.4

0.6

0.8

1.0

300 400 500 600 700 800 900 1000

No

rmaliz

ed

Mo

me

nt

T (K)

0.0 0.1

0.2 0.3

0.4 0.5

0.0

0.2

0.4

0.6

0.8

1.0

300 400 500 600 700 800 900 1000

No

rmaliz

ed

mom

ent

T (K)

0.0 0.1

0.2 0.3

0.4 0.5

127

3.5 Electrical Properties

Electrical properties of prepared samples have been measured using dc-electrical

resistivity and ac-electrical (dielectric) properties measurements system as discussed in

Sec. 2.3.5.

3.5.1 DC-electrical resistivity measurements

Thermal variation of the DC-electrical resistivity (ρ) of Co-Cr substituted series

measured in the temperature range of 298 K to 673 K is shown in Fig. 3.34. The plots of

DC-electrical resistivity (ρ) versus temperature show rapid increase in resistivity with

temperature (metal like behavior), followed by a decrease in resistivity above the metal to

semiconductor transition temperature (TM-S). The metal-semiconductor transition region

for Co-Cr doped samples ranges from 310-380 K. The value of TMS increases slightly

with increase in Co-Cr content up to x ≤ 0.2 after that becomes almost constant. A similar

transition has been reported in Al-Cr substituted Ni-ferrites [252]. Such resistivity-

temperature behavior can be attributed to several factors like occurrence of phase

transition, cation migration, cation re-ordering and presence of impurities. In case of the

samples investigated in this study, it is unlikely that there would be any phase transition

at such low temperature (350-370 K). Also, XRD data did not show the presence of any

impure phase in the materials. Cation migration would also be limited at such a low

temperature and thus cannot be accounted for the observed behavior. A possible and most

likely reason for the observed metallic behavior at ~350-370 K is the presence of

absorbed water in the samples due to the hygroscopic nature of spinel magnesium ferrites

[253]. Although the experiments were performed in dry atmosphere but the possibility of

moisture present deep within the pores of the pellets cannot be ruled out.

Electrical properties of ferrites generally depend upon the composition and

preparation parameters of the synthesized materials [254]. The source of conduction is

the dispersion of the charge carriers of the atoms composing the conductive materials.

The dispersion is produced primarily due to vibration of atoms and/or presence of lattice

defects [255]. Electrical conduction is governed essentially by the manner of generation

of charge carriers and their transportation in a material.

128

The conduction mechanism in ferrites can be explained by the Verwey-de Boer

mechanism [33] in which exchange/hopping of electrons (charge carrier) take place

between ions of the same element with different valence state. Such ions are distributed

randomly over equivalent lattice sites. Hence, conduction in magnesium ferrite originates

mainly from hopping of electron between Fe3+

and Fe2+

ions, when both ions are in

adjacent octahedral sites [37]. The magnitude of this exchange phenomenon depends on

the number of Fe2+

/Fe3+

ion pairs located at the octahedral site and and overlap of orbitals

between transition metal ions and oxygen [36]. Further, the probability of hopping

depends upon the separation between ions involved in conduction process and the

activation energy [34]. The number of such ion pairs depends upon the sintering

conditions and the extent of reduction of Fe3+

to Fe2+

at elevated temperature. In spinel

ferrites, the distance between the cations in A-sites is more than the distance between

cations in B-sites. Thus, the degree of covalency for the A-sites cations is known to be

higher than that of the B-sites cations [256]. Due to these, the mobility of electrons and

holes in the A-sites is expected to be smaller than that in the B-sites.

Figure 3.34: Plot of DC-electrical resistivity of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) versus

temperature; Inset: Arrhenius type plot of ln ρ vs. 103/T for some Co-Cr substituted

magnesium ferrite samples.

129

Room temperature resistivity (ρRT

) of magnesium ferrite samples increases from a

value of 7.50 ×108 Ωcm to 3.47 ×109 Ωcm with the increase in concentration of Co-Cr

binary mixture (Tables 3.18). The value of ρRT

achieved in the present studies is ~102

times greater than those of the order of 106-10

7 Ωcm previously reported [76, 257-260]

for spinel ferrites. So, Co-Cr doped materials can be suitable to curb the eddy current

losses for microwave applications.

Variation in resistivity with dopant concentration is explained based on

occupancy of cations in the sub-lattices of the spinel structure. In Co-Cr doped samples,

substitution of Cr3+

for Fe3+

results in the reduction of Fe3+

ions at B-site (which are in

fact responsible for hopping). Thus, it impedes an electron transfer between the Fe2+

to

Fe3+

ions and consequently increases the resistivity. However, substitution of Co2+

at the

B-site might generate holes by the following exchange reaction, which can also be

involved in conduction:

+ ↔ + (3.4)

The electron hopping however, would mainly contribute in conduction and the overall

effect of doping of binary mixture of Co-Cr is the enhancement of resistivity.

The activation energy of hopping (Ea) is calculated in the temperature range of

373-473 K from a linear plot of lnρ versus 103/T (Inset of Fig. 3.34) in accordance with

the Arrhenius equation [37]. The figure shows that the DC-resistivity decreases linearly

with temperature for all the samples. This can be attributed to the increase in drift

mobility of thermally activated charge carriers. In ferrites, the electrons are localized and

an overlap between the wave functions of ions situated on adjacent sites is small. In the

presence of lattice vibrations, the ions occasionally come so close that the transfer of

electrons from one ion to another occurs with high probability. Hence, the electron

mobility is temperature dependent and is characterized by activation energy. The

observed decrease in DC-resistivity with temperature is normal for semiconductors and

follows Eq. 2.12. The activation energy is determined by fitting the resistivity data to Eq.

2.12. The activation energy of magnesium ferrite is 0.35 eV, which increased to 0.58 eV

with the increase in Co-Cr content up to x = 0.5 (Table 3.18). These values are higher

than the transition energy of Fe2+ and Fe3+ (Ee = 0.2 eV) [261], indicating that the polaron

130

model of electron hopping is favored in the studied samples. The compositional variation

of activation energy is in close agreement with the resistivity data (refer Table 3.18).

The drift mobility (µd) of charge carriers is calculated from the experimental data

of DC-electrical resistivity using Eqs. 2.13-2.14. The magnitude of µd of all the samples

is observed to be in the range of 10-13

-10-14

cm2V

-1s

-1. The drift mobility and resistivity

have an inverse relationship, since samples can have high value electrical resistivity only

because of the low mobility of their charge carriers. The room temperature value of drift

mobility (Table 3.18) shows that it decreases from 34.7×10−14

cm2V

-1s

-1 to 9.6×10

−14

cm2V

-1s

-1, due to electron hopping suppressed at the B-sites causing the increase in

resistivity (Table 3.18) as discussed previously in this section.

Table 3.18: DC-electrical resistivity (ρRT

), activation energy (Ea), drift mobility (µd),

dielectric constant (έ) and dielectric loss tangent (tan δ) of Mg1-xCoxCrxFe2-xO4

(x = 0.0-0.5).


ρRT (Ωcm)10

9 0.75 1.48 2.14 3.02 3.39 3.47

Ea (eV) 0.33 0.41 0.47 0.53 0.57 0.58

µd (cm2V

-1s

-1)10

-14 34.7 18.5 13.5 10.1 9.6 9.9

έ @ 10 kHz 564 457 303 109 98 96

Tanδ @ 10 kHz 0.99 0.66 0.58 0.51 0.51 0.55

The temperature-resistivity plots of Mg1-xMxCrxFe2-xO4 (M = Ni, Cu, Zn and Mn)

are shown in Figs. 3.35-3.38. All the series exhibit semiconductor like behavior with the

increase in temperature and show the variation of resistivity with temperature similar to

that of Co-Cr doped magnesium ferrite samples, as discussed before (Fig. 3.34). The

metal to semiconductor transition temperature (TM-S) ranges from ~320-380 K for all the

synthesized series. The appearance of TM-S is because of adsorbed moisture on the surface

of the sample, as indicated before for Co-Cr doped magnesium ferrite samples.

The variation of DC-electrical resistivity (ρRT

), calculated at the potential

operational range around 300 K, as a function of Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr

contents is shown in Tables 3.19, 3.20, 3.21 and 3.22, respectively. The ρRT

is increased

131

from a value of 7.5×108 to 4.8×10

9 and 7.36×10

9 Ωcm by increasing Ni-Cr and Cu-Cr

contents, respectively (Tables 3.19 and 3.20). Increase in resistivity would cause the

decrease in eddy current losses.

The activation energy of hopping (Ea) for Ni-Cr and Cu-Cr doped samples also

increases from 0.35 to 0.48 eV and 0.60 eV for the Ni-Cr and Cu-Cr substituted Mg-

ferrites, respectively. The observed increase in both the values of ρRT

and Ea for dopant

content can be ascribed to the nature and site occupancy of the cations.

The resistivity of the spinel ferrites is controlled by the Fe2+ / Fe3+ ratio on the

B-sites [262]. Ni2+

and Cu2+

ions have strong tendency to occupy the B-sites [34, 177],

therefore, substitution of nickel and copper replaces Mg2+

ions from the B-sites, having a

negligible contribution to the variation in resistivity. Substitution of Cr3+

ions for Fe3+

at

B-sites decreased the number of iron ions contributing to hopping process, which in turn

leads to the increase in DC-resistivity. These Cr3+ ions do not participate in the

conduction process and also limit the degree of Fe2+

-Fe3+

conduction by blocking up the

Fe2+

-Fe3+

transformation. This phenomenon hinders the Verwey-de-Boer mechanism

[263] between statistically distributed Fe2+

and Fe3+

ions at equivalent lattice sites,

resulting in an increase in resistivity. Furthermore, Cr3+

ions occupy the interstitial B-

sites and results in distortion of the lattice, creating internal stress, which confined the

electron hopping and then reduces the Fe2+

ion generation [264]. Therefore, it is expected

that the resistivity should increase upon doping, as observed. The resistivity of Mg-ferrite

increases almost 5-10 times as a result of substituting the binary mixture of Ni-Cr and

Cu-Cr, as compared to un-doped magnesium ferrite. In addition, Cu-Cr substitution

caused more enhancements in ρRT as compared to Ni-Cr substituted series. This could be

due to valence fluctuations of Fe ions between 2+/3+ valance states and that of Ni

between Ni2+

and Ni3+

that might occur during the annealing process [282, 283]. In this

way, the following exchange interactions may become possible:

Fe3+ + Ni2+ ↔ Fe2+ + Ni3+ (3.5)

As a result of this, the conduction due to hole hopping between Ni2+

/Ni3+

ions

might cause the resistivity to remain suppressed [43]. The room temperature DC-

electrical resistivity (ρRT) also increases with the co-substitution of Zn-Cr and Mn-Cr ions

up to x ≤ 0.3 as shown in Tables 3.21 and 3.22, whereas it start to decrease for further

132

substitution of dopants. The activation energy of hopping (Ea) also varied in a manner

similar to that of ρRT

for the above-said two series. To explain this observed behavior of

both ρRT and Ea, we have to keep in mind again the mechanism owing to the nature and

site occupancy of the cations.

Since Zn2+

prefers the tetrahedral sites [267] and Cr3+

occupies the octahedral

sites, thus co-substitution of Zn-Cr ions would cause the depletion of Mg2+

and Fe3+

at A-

and B-sites, respectively. The rate of electron transfer will slow down with the decrease

of Fe3+ ions which would consequently enhance the DC-resistivity with the increase of

Cr3+

concentration. This happened up to x ≤ 3 beyond which substitution of Zn2+

for Mg2+

causes migration of an equal fraction of Fe3+

from tetrahedral to octahedral sites. This

would result in an increase of Fe3+

concentration at B-sites contributing to the hopping

process and leading to the decrease in DC-resistivity.

In case of Mn-Cr substituted series, an increase in ρRT with the substitution level

of x ≤ 3, is also due to depletion of Fe3+

at B-sites, which slow down the electron transfer

and consequently enhance the DC-resistivity with the increase in Cr3+

concentration.

However, a fall down in the values of ρRT

observed for higher levels of Mn-Cr

substitution (x > 0.3) can be rationalized as, some of the Mn2+

can be oxidized to Mn3+

state during annealing process with an exchange reaction shown below:

Fe3+

+ Mn2+

↔ Fe2+

+ Mn3+

(3.6)

In such a case, both the hole transfer and electron hopping between Mn2+/Mn3+

together with electron hopping between Fe3+

/Fe2+

may also contribute to the conduction

process [268] and therefore decreases the resistivity value for higher levels of

substitution. A similar observation for nano-crystalline ferrites has also been reported by

the other researchers [260]. The variation of drift mobility as a function of Ni-Cr, Cu-Cr,

Zn-Cr and Mn-Cr doped series is presented in Tables 3.19, 3.20, 3.21 and 3.22,

respectively. The µd decreases with dopants contents from a value of 34.7×10−14

cm2V

-1s

-

1 to 6.8×10

−14 cm

2V

-1s

-1 and 4.6×10

−14 cm

2V

-1s

-1 for Ni-Cr (Table 3.19) and Cu-Cr (Table

3.20) substituted samples, respectively. But in case of Zn-Cr and Mn-Cr substituted

series, drift mobility decreases from 34.7×10−14 cm2V-1s-1 to 12.2×10−14 cm2V-1s-1 and

19.6×10−14 cm2V-1s-1 with the increase in dopant contents up to substitution level (x ≤

0.3), while start to increase above this level of substitution (Tables 3.21-3.22). It is

133

obvious that the observed behavior of drift mobility with composition is exactly opposite

to that of the resistivity (Tables 3.19-3.22). This is rationalized that the samples may have

a high electrical resistivity only because of the low mobility of their charge carriers and

vice versa. The data for the drift mobility is in concordance with the electrical resistivity

and the activation energy.

The electrical properties of nanosized Mg-ferrite change drastically with variation

of crystallite size from bulk to the nanoscale. The room temperature electrical resistivity

for bulk ferrite (191 nm) is reported as 1.5×108 Ωcm [269], while higher values of

resistivity has been obtained [34, 37] as 7.5 ×108 Ωcm for magnesium ferrite

nanomaterials having crystallite size of 47 nm. This suggests that the nanomaterials are

more appropriate for microwave applications, relative to their bulk counterparts; as the

higher resistivity is pre-requisite for microwave applications to curb the eddy current

losses, which is important for the ferrite performance [270].

The greater value of resistivity in Mg-ferrite nanoparticles might be dependent on

the grain size. The increase in resistivity with grain size reduction, therefore, can be

attributed to the size effects as reported in the literature [271, 272] and also due to the

increase in grain boundary volume and the associated impedance to the flow of charge

carriers. It is known that the smaller grains with prominent grain boundary lead to a

higher resistivity. Therefore, ferrites with high resistivity have negligible bulk eddy

current loss. As there is a large number of grain boundary regions and a smaller contact

between the adjacent grains, hence, the presence of grain boundaries acted as an

impediment to the domain wall motion. Prominent grain boundaries also implied smaller

grains and a higher resistivity or there was a reduced electron flow. The grain boundaries

are insulating in nature and are highly resistive to the flow of electrons in the material.

Ferrites are composed of two layers, first is conducting layer consisting of a large number

of grains and second layer consist of grain boundaries with insulating nature. The

samples with low porosity and small grain size have a large number of grain boundaries

and would be highly resistant. The resistance of the grain boundaries in the material

arises due to the development of potential barrier at the surface of grain boundaries. This

potential barrier is developed due to trapping of charge carriers (electrons in ferrites) at

the boundary and the compensating opposite space charges inside the grain (holes in case

134

of ferrites). This increase of the number of grain boundaries in a sample result in the

increases of the resistivity of the samples [273]. Apart from this intrinsic resistance to the

flow of electrons, another explanation for this phenomenon is that the presence of grain

boundaries in the samples result in the bond breakage which also affects the electron

exchange between Fe+2

and Fe+3

[274]. The high resistive samples may have reduced

eddy current losses, thus, the substituted magnesium ferrite samples in the present study

can be suitable for microwave applications.

It is concluded that the synthesized samples show metal to semiconductor

transition in the temperature range of 310-380 K and thereafter obey the Arrhenius-type

temperature dependence which shows a behavior similar to that of semiconductors. The

DC-resistivity of substituted Mg-ferrite continued to increase by doping with binary

mixtures of Ni-Cr and Cu-Cr ions while it committed to decrease after specific level of

substitution (x > 0.3) with Zn-Cr and Mn-Cr ions. The observed behavior of drift

mobility with the dopant contents is exactly opposite to that of DC-resistivity. The higher

electrical resistivity of samples is in concordance with their high activation energy of

hopping and low drift mobility.

Figure 3.35: Plot of DC-electrical resistivity of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5) versus

temperature. Inset: Arrhenius type plot of ln ρ vs. 103/T for some Ni-Cr substituted


135

Figure 3.36: Plot of DC-electrical resistivity of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5) versus

temperature; Inset: Arrhenius type plot of ln ρ vs. 103/T for some Cu-Cr substituted


Figure 3.37: Plot of DC-electrical resistivity of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) versus

temperature. Inset: Arrhenius type plot of ln ρ vs. 103/T for some Zn-Cr substituted


136

Figure 3.38: Plot of DC-electrical resistivity of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5) versus

temperature. Inset: Arrhenius type plot of ln ρ vs. 103/T for some Mn-Cr substituted




dielectric constant (έ) and dielectric loss tangent (tanδ) of Mg1-xNixCrxFe2-xO4

(x = 0.0-0.5).


ρRT (Ωcm)10

9 0.75 1.59 2.93 4.29 4.85 4.31

Ea (eV) 0.33 0.38 0.41 0.45 0.48 0.46

µd (cm2V

-1s

-1)10

-14 34.7 17.4 9.9 7.2 6.8 8.2

έ @ 10 kHz 564 364 242 135 81 57

Tanδ @ 10 kHz 0.99 1.28 1.05 0.99 0.87 1.72

137



dielectric constant (έ) and dielectric loss tangent (tan δ) of Mg1-xCuxCrxFe2-xO4

(x = 0.0-0.5).


ρRT (Ωcm)10

9 0.75 1.32 2.02 3.72 5.20 7.36

Ea (eV) 0.33 0.37 0.43 0.51 0.56 0.60

µd (cm2V

-1s

-1)10

-14 34.7 20.5 14.1 8.1 6.1 4.6

έ @ 10 kHz 564 396 293 89 76 42

Tanδ @ 10 kHz 0.99 1.73 1.39 1.18 2.45 2.97

Table 3.21: DC-electrical resistivity (ρRT), activation energy (Ea), drift mobility (µd),

dielectric constant (έ) and dielectric loss tangent (tan δ) of Mg1-xZnxCrxFe2-xO4

(x = 0.0-0.5).


ρRT (Ωcm)10

9 0.75 1.62 1.97 2.56 1.64 1.17

Ea (eV) 0.33 0.39 0.44 0.46 0.40 0.36

µd (cm2V

-1s

-1)10

-14 34.7 17.1 15.1 12.2 20.0 29.6

έ @ 10 kHz 564 378 273 121 182 208

Tanδ @ 10 kHz 0.99 0.76 0.68 0.49 0.64 0.75

138



dielectric constant (έ) and dielectric loss tangent (tan δ) of Mg1-xMnxCrxFe2-xO4

(x = 0.0-0.5).


ρRT (Ωcm)10

9 0.75 0.92 1.18 1.56 0.55 0.41

Ea (eV) 0.33 0.34 0.37 0.40 0.32 0.29

µd (cm2V

-1s

-1)10

-14 34.7 29.8 24.6 19.6 58.8 84.1

έ @ 10 kHz 564 425 375 222 515 589

Tanδ @ 10 kHz 0.99 0.86 0.78 1.51 2.93 3.55

3.5.2 Dielectric measurements

The dielectric measurements reveal an insight into the behavior of electrical

charge carriers. The dielectric strength of any material essentially exhibits the inherent

ability of the material to withstand a voltage applied without undergoing any structural

degradation or showing conduction. Ferrites are very resistant materials, therefore

exhibits low dielectric strengths. Dielectric constants of prepared samples are calculated

using Eq. 2.15.

The change in the real part of the dielectric constant (έ) and dielectric loss tangent

(tan δ) for Co-Cr doped magnesium ferrites with applied frequency ranging from 100 Hz

to 3 MHz are shown in Figs. 3.39 and 3.40, respectively. It is found that έ and tanδ are

normally quite high but decreases rapidly with an increase in frequency in the low

frequency regime. However, the dispersion in both the έ and tanδ slows down at higher

frequency, which is normal behavior of ferrite materials. A similar variation of the

dielectric constant with frequency was observed in the cases of ZnFe2O4, MnFe2O4 [275]

and ZnxCu1-xFe2O4, MnxCu1-xFe2O4 [276].

139

Figure 3.39: Dielectric constant (έ) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) versus ln f.

Figure 3.40: Dielectric loss tangent (tanδ) of Mg1-xCoxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f.

0

4

8

12

16

20

4 6 8 10 12 14 16

έ(1

03)

ln f

0.0 0.1

0.2 0.3

0.4 0.5

0

4

8

12

16

20

24

4 6 8 10 12 14 16

tanδ

ln f

0.0 0.1

0.2 0.3

0.4 0.5

The dielectric properties of heterogeneous systems like ferrites are dependent on

several factors including the method of preparation, sintering temperature, sintering

atmosphere, etc. Since, ferrite powdered samples are sintered under

conditions; the divalent iron Fe

material leading to high conductivity grains. When such a material is cooled in an

atmosphere of oxygen, it is possible to form layers of very low conductivity over its

constituent grains. Almost, all the ferrites in the polycrystalline form have such high

conductivity grains separated by low conductivity layers so that they behave as

heterogeneous dielectric materials

dielectric constant as high as 10

The dielectric behavior

interfacial polarization

frequency, the dipolar

dielectric constant but at

significant. A physical interpretation of dispersion can be

hopping between Fe2+

occurs through a mechanism

the electronic exchange, as Fe

is observed in the direction of applied electric field,

At low frequency, the electron exchange between Fe

the alternating field but the frequency of electron exchange between two ions cannot

follow the high alternating field. This results in a decrease in polarization and the

conduction lag beyond a certain frequency of

value of dielectric constant is lower at high frequencies than at low frequencies.

The variation of dielectric constant and dielectric loss tangent with the Co

contents is shown in Table 3.18

a frequency of 10 kHz, for different

0.0) to 96 (x = 0.5). These values are 10

magnesium ferrite. El-

(at 1 kHz) for the bulk MgFe

101–10

2 for the Co-Cr doped samples at the same temperature and 10 kHz, which are two



atmosphere, etc. Since, ferrite powdered samples are sintered under

the divalent iron Fe2+

is expected to be formed in the body of the ferrite



nstituent grains. Almost, all the ferrites in the polycrystalline form have such high


heterogeneous dielectric materials [277]. Due to the heterogeneous dielectric behavior,

electric constant as high as 103-10

5 is observed in the case of ferrites at low frequencies.

The dielectric behavior of the spinel ferrites can be explained by considering

interfacial polarization at the interfaces of grains and grain boundaries

dipolar and interfacial polarization both contribute to the value of

dielectric constant but at higher frequencies, only the electronic polarization becomes

A physical interpretation of dispersion can be made on the bas

2+ and Fe

3+ ions, as the dielectric polarization

through a mechanism similar to that of electrical conduction process

electronic exchange, as Fe2+ Fe3++e−1, the local displacement

in the direction of applied electric field, enable to determine the polarization.

At low frequency, the electron exchange between Fe2+

and Fe3+

ions is capable to follow



conduction lag beyond a certain frequency of an externally applied field. Therefore, the


The variation of dielectric constant and dielectric loss tangent with the Co

Table 3.18. The room temperature value of dielectric constant (

frequency of 10 kHz, for different Co-Cr contents is observed to decrease from 564 (x =

0.0) to 96 (x = 0.5). These values are 102 to 103 times smaller than thos

-Hiti [280] has obtained a value on the order of

(at 1 kHz) for the bulk MgFe2O4. In the present study, έ is found to be in the order of

Cr doped samples at the same temperature and 10 kHz, which are two

140



atmosphere, etc. Since, ferrite powdered samples are sintered under slightly reducing

is expected to be formed in the body of the ferrite



nstituent grains. Almost, all the ferrites in the polycrystalline form have such high


. Due to the heterogeneous dielectric behavior,

is observed in the case of ferrites at low frequencies.

of the spinel ferrites can be explained by considering the

of grains and grain boundaries [278]. At low

polarization both contribute to the value of

, only the electronic polarization becomes

on the basis of electron

dielectric polarization process in ferrites

conduction process [279]. Due to

nt of the charge carriers

determine the polarization.

ions is capable to follow



applied field. Therefore, the


The variation of dielectric constant and dielectric loss tangent with the Co-Cr

dielectric constant (έ) at

is observed to decrease from 564 (x =

times smaller than those of bulk

n the order of 104 for έ at 312 K

is found to be in the order of

Cr doped samples at the same temperature and 10 kHz, which are two

141

orders of magnitude smaller than that of the bulk MgFe2O4. The considerably higher

values at low frequency can be attributed to higher capacitance of the grain boundary. It

can be seen from Fig. 3.39 that the dispersion of έ with the frequency is the maximum for

undoped Mg-ferrite and magnitude of dispersion decreases as more Co-Cr is

incorporated. The decrease in the value of έ with Co-Cr contents is due to reduction in Fe

contents to be polarized (space charge carriers) at the B-site, resulting in a low space

charge polarization governed by space charge carriers and resistivity of the materials.

The value of dielectric loss tangent (tan δ) decreases from 0.99 to 0.51 at 10 kHz

for Co-Cr contents up to x = 0.4 (Table 3.18). Iwauchi [281] described that the dielectric

losses in ferrites are usually a reflection of electrical conductivity; thus, the materials with

high electrical resistivity exhibit low dielectric constant and loss tangent. Hence,

electrical resistivity and dielectric properties both are transport properties whose

variations are proportional to the sample composition and most probably the mechanism

responsible for these two phenomena may be the same or similar to each other. Data of

the dielectric parameters support the data of electrical resistivity measurements to much

extent (Table 3.18).

The variation of dielectric constant and dielectric loss tangent for Ni-Cr, Cu-Cr,

Zn-Cr and Mn-Cr doped magnesium ferrites with frequency are shown in Figs. 3.41-3.48.

It is obvious from these plots that the dielectric parameters (έ & tan δ) show dispersion

phenomenon similar to that of the Co-Cr doped samples, which is a usual dielectric

behavior, except the dielectric loss tangent of Cu-Cr substituted samples exhibiting the

loss peak. The relaxation peak occurs when the hopping frequency of electron from Fe2+

to Fe3+ coincide with the applied field frequency. In Cu-Cr doped samples, hopping

frequency range of 300 Hz to 2 kHz for charge carriers coincides with the applied field

frequency leading to a relaxation peak.

The Koop’s theory [277] and Maxwell-Wagner bi-layer model [218] are ideal to

explain the dispersion in dielectric materials (as discussed above). Ferrite materials are

composed of well conducting grains separated by insulating grain boundaries, which

cause hindrance in the conduction process. If grain boundaries possess the larger

resistance, then the charge carriers get aligned at the grain boundaries and results into the

polarization of dielectric medium which leads to a greater dielectric constant [265]. As,

142

the frequency increases, the increase in electrical resistivity due to dopant substitution

reduces the probability of charge carriers to approach the grain boundaries. This distorts

the polarization setup in material resulting in the decrease of dielectric constant at higher

frequencies [282].

In the low frequency region, which corresponds to high resistivity (due to grain

boundaries), more energy is required for electron exchange between Fe2+

and Fe3+

ions

and thus the energy loss is high. In the high frequency range, which corresponds to low

resistivity (due to the grains), a small energy is needed for electron transfer between Fe2+

and Fe3+

ions in the grains and accompanied by a small eddy current and hence a

decrease in the energy loss. Consequently, the values of dielectric constant and loss

tangent decrease substantially with increasing frequency.

Figure 3.41: Dielectric constant (έ) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5) versus ln f

0

4

8

12

16

20

4 6 8 10 12 14 16

έ (

10

3)

ln f

0.0 0.1

0.2 0.3

0.4 0.5

143

Figure 3.42: Dielectric loss tangent (tanδ) of Mg1-xNixCrxFe2-xO4 (x = 0.0-0.5) versus

ln f.

Figure 3.43: Dielectric constant (έ) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5) versus ln f.

0

4

8

12

16

20

24

4 6 8 10 12 14 16

tanδ

ln f

0.0 0.1

0.2 0.3

0.4 0.5

0

4

8

12

16

20

4 6 8 10 12 14 16

έ (

10

3)

ln f

0.0 0.1

0.2 0.3

0.4 0.5

144

Figure 3.44: Dielectric loss tangent (tanδ) of Mg1-xCuxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f.

Figure 3.45: Dielectric constant (έ) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) versus ln f.

0

4

8

12

16

20

24

4 6 8 10 12 14 16

tanδ

ln f

0.0 0.1

0.2 0.3

0.4 0.5

0

4

8

12

16

20

4 6 8 10 12 14 16

έ(1

03)

ln f

0.0 0.1

0.2 0.3

0.4 0.5

145

Figure 3.46: Dielectric loss tangent (tanδ) of Mg1-xZnxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f.

Figure 3.47: Dielectric constant (έ) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5) versus ln f.

0

4

8

12

16

20

24

4 6 8 10 12 14 16

tanδ

ln f

0.0 0.1

0.2 0.3

0.4 0.5

0

5

10

15

20

25

30

35

4 6 8 10 12 14 16

έ (

10

3)

ln f

0.0 0.1

0.2 0.3

0.4 0.5

146

Figure 3.48: Dielectric loss tangent (tanδ) of Mg1-xMnxCrxFe2-xO4 (x = 0.0-0.5) versus

ln f.

The compositional dependence of dielectric constant and dielectric loss tangent at

10 kHz for Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr substituted magnesium ferrites are presented

in Tables 3.19, 3.20, 3.21 and 3.22, respectively. Introducing Ni-Cr and Cu-Cr into the

magnesium ferrite, έ decreases to 57 and 42, while tanδ increases to 1.72 and 2.97 at 10

kHz, respectively, for substitution level up to x = 0.5 (Tables 3.19 and 3.20). The έ value

decreases to 121 and 222 at a frequency of 10 kHz for the Zn-Cr and Mn-Cr substituted

magnesium ferrites respectively, as shown in Tables 3.21-3.22. The value of tanδ

decreases to 0.49, while increases to 3.55 for Zn-Cr and Mn-Cr substituted magnesium

ferrites respectively, at the same frequency. The formation of an insulating inter-granular

layer embodiment of grains of the samples reduces the oxidation rate of Fe2+

ions results

in the decline of dielectric parameters. El-Hiti [280] has obtained a tan δ value of the

order 101 at 312 K (at 1 kHz) for the bulk MgFe2O4 prepared by ceramic method. In the

present study, tan δ for all samples is about 0.5-3.55 at 300 K and 10 kHz. The dielectric

loss of Mg-ferrite and its doped derivatives is thus found to be lower compared to that of

their bulk counterparts [280].

0

4

8

12

16

20

24

28

32

4 6 8 10 12 14 16

tanδ

ln f

0.0 0.1

0.2 0.3

0.4 0.5

147

It has been found that the dielectric parameters behave in a manner opposite to

that of the resistivity and the activation energy so that the dielectric polarization

mechanism is similar to that of electrical conduction which is in close agreement with the

Iwauchi’s assumption [275]. In other words, an increase in electrical resistivity caused

the decrease in dielectric parameters of magnesium ferrite by doping it with different

contents of Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr ions. The high electrical resistivity along

with low dielectric constant and dielectric losses are expected to fulfill the requirement of

microwave applications.

Thus, it can be concluded that the dielectric constant (έ) and dielectric loss tangent

(tan δ) of magnesium ferrite and its derivatives show dispersion with increasing applied

frequency. The maximum dispersion is observed at low frequency while it is less at

higher frequency, obeying the Maxwell-Wagner model for dielectric materials. The

dielectric parameters (έ, tanδ) of magnesium ferrite can be reduced by doping with M-Cr

(M = Co, Ni, Cu, Zn and Mn) ions.

Conclusions

An attempt is made to meet the challenges for the advancements in the new ferrite

technology. The hypothetical variation in the structural, magnetic and electrical

properties of cubic spinel magnesium ferrites introduced by the substitution of binary

mixtures of cations have been rationalized and proved. The polyethylene glycol (PEG)

assisted microemulsion method is a reliable and economic method to synthesize

nanoparticles of magnesium ferrites (MgFe2O4) doped with the binary mixtures of Co-Cr,

Ni-Cr, Cu-Cr, Zn-Cr and Mn-Cr ions and characterized by utilizing various advanced

techniques. The obtained results are summarized with following concluding remarks:

• XRD analysis revealed that all the samples have single phase cubic spinel

structure except Cu-Cr doped samples which possessed slight tetragonal distortion

in the cubic structure which might be due to Cu2+ ion substitution in MgFe2O4

• Crystallite size lied below 50 nm except for higher levels of Cu-Cr doped samples

and small enough to obtain the suitable signals-to-noise ratio for switching

applications

148

• All the samples are slightly porous evident from SEM analysis and supported by

comparing the X-ray density and bulk density data

• The cation distribution revealed by the Mössbauer analysis are found helpful to

elaborate the data of magnetic and electrical measurements carried out here

• It is observed that the saturation magnetization, the remanence, the coercivity and

the magnetocrystalline anisotropy coefficient (K1) increased with decreasing

temperature, while varied in a different manner with the dopant contents for each

series of doped Mg-ferrites

• Co-Cr doped samples have significant enhancement in MS value among all the

synthesized series, particular sample with composition x = 0.3 possesses the

highest value of saturation magnetization

• Coercivity has more significant increase for Co-Cr co-dopant with the decrease in

temperature which might be due to more anisotropic nature of Co ions

• Ni-Cr doped samples have the highest value of TC among all the synthesized

series and particularly sample with composition x = 0.5 has the greater value of

Curie temperature

• DC-resistivity of the order of 109

for doped Mg-ferrite has been increased almost

100 times as compared to that reported previously for spinel ferrites

• Cu-Cr doped samples have the highest value of ρRT among all the synthesized

series and sample with composition x = 0.5 has the highest value in particular

• The dielectric constant and dielectric loss decreases with the increase in frequency

in all the samples of the five series investigated here

• Dielectric constant and dielectric loss tangent decreased with the increase in

dopant contents except for the Mn-Cr and Zn-Cr doped Mg-ferrites

• The observed variations of all these magnetic and electrical properties have been

interpreted in terms of the effects of co-substituents on site occupancies of the

cations and the law of approach to saturation has been applied successfully and

calculated values of MS and K1 have a high accuracy

• Co-substitution of M-Cr (Co, Ni, Cu, Zn and Mn) ions has brought remarkable

changes in the temperature dependence of the magnetization, coercivity, magnetic

149

anisotropy and electrical properties of polycrystalline magnesium ferrite i.e. Mg1-

xMxCrxFe2-xO4 (x = 0.0-0.5).

• The synthesized nanomaterials possess high saturation magnetization, low

coercivity and improved room temperature resistivity together with low dielectric

loss

• Data obtained demonstrate the ability to tune properties of doped Mg-ferrite to

match intended applications.

• Samples with compositions, Co-Cr content (x) = 0.3, Ni-Cr content (x) = 0.5 and

Cu-Cr content (x) = 0.5 can be suitable for some practical applications

Suggestions for Further Research

Magnetic properties have been studied at normal and low temperature; these

properties can be performed at high temperature to see the detailed effects of temperature

on various parameters of doped magnesium ferrites.

Magnetoelastic properties i.e. magnetostrictive and megnetoresistive properties of

spinel Mg-ferrite can be performed for torque/stress sensor and actuator applications.

150

References

[1] Safarik, I.; Safarikova, M. Magnetic Nanoparticles and Biosciences, Monotshefte

fur Chemie, Frankfurt, 2002, 133, 737-759.

[2] Buzea, C.; Pacheco, I.; Robbie, K. Biointerphases, 2007, 2, MR17- MR71.

[3] Allhoff, F.; Lin, p.; Moore, D. What is Nanotechnology and why does it Matter?

John Wiley and Sons, New York, 2010, pp. 3.

[4] Sohn, B.H.; Cohen, R.E. Chem. Mater. 1997, 9, 264-269.

[5] Klabunde, K.J. Nanoscale Materials in Chemistry, Wiley, New York, 2001, pp.1.

[6] Adams, D.M. Inorganic Solids, John Willey, London, 1974, pp. 68.

[7] Smyth, D.M. The Defect Chemistry of Metal Oxides, Oxford University Press,

New York, 2000, Chap. 2, pp. 35.

[8] West, A.R. Solid State Chemistry and its Applications, John Wiley & Sons,

Singapore, 1989, Chap.16, pp. 571.

[9] Sickafus, K.E.; Wills, J. J. Am. Ceram. Soc. 1999, 82, 3279-3281.

[10] Smit, J.; Wijn, H.P.J. Ferrites-Physical Properties of Ferrimagnetic Oxides in

Relation to their Technical Applications, John Wiley and Sons, New York,

1959, pp. 369.

[11] Goldmann, A. Modern Ferrite Technology, Van Nosttrand Reinhold, New York,

1990, pp. 71.

[12] Craik, D.J. Magnetic Oxide, Part 1, John Wiley & Sons, Bristol, 1975, ch. 9, part

2.

[13] Verway, E.J.W.; Heilmann, E.L. J. Chem. Phys. 1947, 15, 174-180.

[14] Romeign, F.C. Philips Res. Rep. 1953, 8, 304-342.

[15] Franco, A.; Silva, M.S. J. Appl. Phys. 2011, 109, 07B505.

[16] Seshan, K.; Shashimohan, A.L.; Chakrabarty, D.K.; Biswas, A.B. Phys. Status

Solidi A, 1981, 68, 97-101.

[17] Bahout, M.M.; Bertrand, S.; Pena, O. J. Solid State Chem. 2005, 178, 1080-1086.

[18] Mathew, D.S.; Juang, R.S. Chem. Eng. J. 2007, 129, 51-65.

[19] King, R.J. Geology Today, 2004, 20, 194-200.

151

[20] Deer, W.A.; Howei, R.A.; Zussman, J. An Introduction in Rock-Forming

Minerals, 2nd Ed, Longmann, London, 1992, pp. 58.

[21] Modi, K.B.; Joshi, H.H.; Kulkarni, R.G. J. Mater. Sci. 1996, 31, 1311-1317.

[22] Chen, Q.; Zhang, Z.J. Appl. Phys. Lett. 1998, 73, 3156-3158.

[23] Maensiri, S.; Sangmanee, M.; Wiengmoon, A. Nanoscale Res. Lett. 2009, 4, 221-

228.

[24] Hankare, P.P.; Vader, V.T.; Patil, N.M.; Jadhav, S.D.; Sankpal, U.B.; Kadam,

M.R.; Chougule, B.K.; Gajbhiye, N.S. Mater. Chem. Phys. 2009, 113, 233-238.

[25] Chen, C. Magnetism and Metallurgy of Soft Magnetic Materials, Dover

Publications, New York, 1986, pp. 26

[26] Zhao, L.; Zhang, H.; Xing, Y.; Song, S.; Yu, S.; Shi, W.; Guo, X.; Yang, J.; Lei,

Y.; Cao, F. J. Solid State Chem. 2008, 181, 245-252.

[27] Morrish, A.H. Physical Principles of Magnetism, John Wiley & Sons, New York,

1965, pp. 487.

[28] Pelecky, D.L.; Rieke, R.D. Chem. Mater. 1996, 8, 1770-1783.

[29] Néel, L. Ann. Phys. (Paris), 1948, 3, 137-198.

[30] Goldman, A. Modern Ferrite Technology, 2nd

Ed. Springer Science Inc. New

York, 2006, Ch. 8, pp. 217.

[31] Wang, L.; Zhang, Q. J. Alloys Compd. 2009, 469, 251-257.

[32] West, A.R. Solid State Chemistry and its Applications; John Wiley & Sons,

Singapore, 2003.

[33] Verwey, E.J.W.; Haayman, P.W.; Romeijn, C.W. J. Chem. Phys. 1947, 15,

181-187.

[34] Iqbal, M.J.; Ahmad, Z.; Meydan, T.; Nlebedim, I.C. Mater. Res. Bull. 2012, 47,

344-351.

[35] Almeida, R.M.; Paraguassu, W.; Pires, D.S.; Corrêa, R.R. Paschoal, 2009, 35,

2443-2447.

[36] Iqbal, M.J.; Ashiq, M.N.; Gomez, P.H. J. Magn. Magn. Mater. 2008, 320, 881-

886.

152

[37] Iqbal, M.J.; Ahmad, Z.; Meydan, T.; Melikhov, Y. J. App. Phys. 2012, 111,

033906.

[38] El-Hiti, M.A.; El Ata, A.M.A. J. Magn. Magn. Mater. 1999, 195, 667-678.

[39] Kittel, C. Introduction to Solid State Physics, John Wiley and Sons, New York,

Ch. 11, pp. 389.

[40] Ikuta, H.; Takanaka, K.; Wakihara, M. Thermochim. Acta, 2004, 414, 227-232.

[41] Kao, K.C. Dielectric Phenomena in Solids, Elsevier Academic Press, London,

2004, pp. 41.

[42] Pillai, O.S. Solid State Physics, 5th

Ed. New Age International, New Delhi, 2003,

pp. 760.

[43] Lixi, W.; Qiang, H.; Lei, M.; Qitu, Z. J. Rare Earth, 2007, 25, 216-219.

[44] Hyeon, T.; Chung, Y.; Park, J.; Lee, S.S.; Kim, Y.W.; Park, B.H. J. Phys. Chem.

B 2002, 106, 6831-6833.

[45] Kong, L.B.; Li, Z.W.; Lin, G.Q.; Gan, Y.B. J. Am. Ceram. Soc. 2007, 90, 2104-

2112.

[46] Koseoglu, Y.; Kavas, H.; Aktas, B. Phys. Stat. Sol. (a) 2006, 203, 1595-1601.

[47] Willey, R.J.; Noirclerc, P.; Busca, G. Chem. Eng. Commun. 1993, 123, 1-16.

[48] Shimizu, Y.; Arai, H.; Seiyama, T. Sens. Actuators, 1985, 7, 11-22.

[49] Chen, N.S.; Yang, X.J.; Liu, E.S.; Huang, J.L. Sens. Actuators B, 2000, 66,

178-180.

[50] Konishi, K.; Maehara, T.; Kamimori, T.; Aono, H.; Naohara, T.; Kikkawa, H.;

Watanabe, Y.; Kawachi, K. J. Magn. Magn. Mater. 2004, 272-276, 2428-2429.

[51] Haart, L.G.J.; Blasse, G. Solid State Ionics, 1985, 16, 137-139.

[52] Huang, Y.; Tang, Y.; Wang, J.; Chen, Q. Mater. Chem. Phys. 2006, 97, 394-397.

[53] Coquay, P.; Peigney, A.; De Grave, E.; Vandenberghe, R.E.; Laurent, C. J. Phys.

Chem. B, 2002, 106, 13199-13210.

[54] Astorino, E.; Busca, G.; Ramis, G.; Willey, R.J. Catal. Lett. 1994, 23, 353-360.

[55] Busca, G.; Finocchio, E.; Lorenzelli, V.; Trombetta, M.; Rossini, S.A. J. Chem.

Soc. Faraday Trans. 1996, 92, 4687-4693.

153

[56] Sheongi, S.W.; Mostovoy, M. Nat. Mater. 2007, 6, 13-20.

[57] Singh, A.; Pandey, V.; Kotnala, R.K.; Pandey, D. Phys. Rev. Lett. 2008, 101,

247602.

[58] Bahadur, D. Bull. Mater. Sci. 1992, 15, 431-439.

[59] Oliver, S.A.; Willey, R.J.; Hamdeh, H.H.; Oliveri, G.; Busca, G. Scr. Metall.

Mater. 1995, 33, 1695-1701.

[60] Bhosale, D.N.; Choudhari, N.D.; Sawant, S.R.; Bakare, P.P. J. Magn. Magn.

Mater. 1997, 173, 51-58.

[61] Zimnol, M.; Graff, A.; Sieber, H.; Senz, S.; Schmidt, S.; Mattheis, R.; Hesse, D.

Solid State Ionics, 1997, 101-103, 667-672.

[62] Kuznetsov, M.V.; Pankhurst, Q.A.; Parkin, I.P. J. Mater. Chem. 1998, 8, 2701-

2706.

[63] Chandrasekaran, G.; Sebastian, P.N. Mater. Lett. 1998, 37, 17-20.

[64] Chen, Q.; Rondinone, A.J.; Chakoumakos, B.C.; Zhang, Z.J. J. Magn. Magn.

Mater. 1999, 194, 1-7.

[65] Ravinder, D.; Latha, K. Mater. Lett. 1999, 41, 247-253.

[66] Liu, C.; Zou, B.; Rondinone, A.J.; Zhang, Z.J. J. Am. Chem. Soc. 2000, 122,

6263-6267.

[67] Sěpelak, V.; Schultze, D.; Krumeich, F.; Steinike, U.; Becker, K.D. Solid State

Ionics, 2001, 141, 677-682.

[68] Men, A.N.; Fetisov, V.B.; Vorobiev, Y.P.; Mesnyankina, S.L.; Kamyshov, V.M.

J. Phys. Chem. Solids, 2001, 62, 635-640.

[69] Rana, M.U.; Abbas, T. Mater. Lett. 2002, 57, 925-928.

[70] Qi, X.; Zhou, J.; Yue, Z.; Gui, Z.; Li, L. J. Magn. Magn. Mater. 2002, 251, 316-

322.

154

[71] Antoshina, L.G.; Goryaga, A.N.; Kokorev, A.I. J. Magn. Magn. Mater. 2003,

258-259, 516-519.

[72] Sěpelak, V.; Baabe, D.; Mienert, D.; Litterst, F.J.; Becke, K.D. Scr. Mater. 2003,

48, 961-966.

[73] Radwan, F.A.; Ahmed, M.A.; Abdelatif, G. J. Phys. Chem. Solids, 2003, 64,

2465-2477.

[74] Rabanal, M.E.; Várez, A.; Levenfeld, B.; Torralba, J.M. J. Mater. Process.

Technol. 2003, 143-144, 470-474.

[75] Chauhan, B.S.; Kumar, R.; Jadhav, K.M.; Singh, M. J. Magn. Magn. Mater.

2004, 283, 71-81.

[76] Chhaya, S.D.; Pandya, M.P.; Chhantbar, M.C.; Modi, K.B.; Baldha, G.J.; Jo, H.H.

J. Alloys Compd. 2004, 377, 155-161.

[77] Turkin, A.I.; Drebushchak, V.A. J. Cryst. Growth, 2004, 265, 165-167.

[78] Verma, S.; Joy, P.A.; Khollam, Y.B.; Potdar, H.S.; Deshpande, S.B. Mater. Lett.

2004, 58, 1092-1095.

[79] Thummer, K.P.; Chhantbar, M.C.; Modi, K.B.; Baldha, G.J.; Joshi, H.H. Mater.

Lett. 2004, 58, 2248-2251.

[80] Liu, Y.; Liu, Z.; Yang, Y.; Yang, H.; Shen, G.; Yu, R. Sens. Actuators, B 2005,

107, 600-604.

[81] Pradeep, A.; Chandrasekaran, G. Mater. Lett. 2006, 60, 371-374.

[82] Lakshman, A.; Rao, P.S.V.S.; Rao, B.P.; Rao, K.H. J. Phys. D: Appl. Phys.

2005, 38, 673-678.

[83] Masti, S.A.; Sharma, A.K.; Vasambekar, P.N.; Vaingankar, A.S. J. Magn. Magn.

Mater. 2006, 305, 436-439.

[84] Modi, K.B.; Chhantbar, M.C.; Joshi, H.H. Ceram. Int. 2006, 32, 111-114.

155

[85] Bergmann, I.; Šepelák, V.; Becker, K.D. Solid State Ionics, 2006, 177, 1865-

1868.

[86] Lakshman, A.; Rao, P.S.V.S.; Rao, K.H. Mater. Lett. 2006, 60, 7-10.

[87] Kong, L.B.; Li, Z.W.; Lin, G.Q.; Gan, Y.B. J. Am. Ceram. Soc. 2007, 90, 3106-

3112.

[88] Šepelák, V.; Bergmann, I.; Menzel, D.; Feldhoff, A.; Heitjan, P.; Litterst, F.J.;

Becker, K.D. J. Magn. Magn. Mater. 2007, 316, e764-e767.

[89] Ichiyanagi, Y.; Kubota, M.; Moritake, S.; Kanazawa, Y.; Yamada, T.; Uehashi, T.

J. Magn. Magn. Mater. 2007, 310, 2378-2380.

[90] Deng, H.; Chen, H.; Li, H. Mater. Chem. Phys. 2007, 101, 509-513.

[91] Vital, A.; Angermann, A.; Dittmann, R.; Graule, T.; Topfer, J. Acta Mater. 2007,

55, 1955-1964.

[92] Shah, J.; Kotnala, R.K.; Singh, B.; Kishan, H. Sens. Actuators, B 2007, 128, 306-

311.

[93] Sharma, S.K.; Kumar, R.; Kumar, S.; Kumar, V.V.S.; Knobel, M.; Reddy, V.R.;

Banerjee, A.; Singh, M. Solid State Commun. 2007, 141, 203-208.

[94] Pradeep, A.; Priyadharsini, P.; Chandrasekaran, G. J. Magn. Magn. Mater. 2008,

320, 2774-2779.

[95] Li, S.; Wang, E.; Tian, C.; Mao, B.; Kang, Z.; Li, Q.; Sun, G. J. Solid State Chem.

2008, 181, 1650-1658.

[96] Barati, M.R. J. Alloys Compd. 2009, 478, 375-380.

[97] Gadkari, A.B.; Shinde, T.J.; Vasambekar, P.N. Mater. Charact. 2009, 60, 1328-

1333.

[98] Gaudon, M.; Pailhe, N.; Wattiaux, A.; Demourgues, A. Mater. Res. Bull. 2009,

44, 479-484.

156

[99] Hankare, P.P.; Vader, V.T.; Sankpal, U.B.; Gavali, L.V.; Sasikala, R.; Mulla, I.S.

Solid State Sci. 2009, 11, 2075-2079.

[100] Chand, J.; Singh, M. J. Alloys Compd. 2009, 486, 376-379.

[101] Sasaki, T.; Ohara, S.; Naka, T.; Vejpravova, J.; Sechovsky, V.; Umetsu, M.;

Takami, S.; Jeyadevan, B.; Adschiri, T. J. Supercrit. Fluids, 2010, 53, 92-94.

[102] Ahmed, Y.M.Z.; Ewais, E.M.M.; Zaki, Z.I. J. Alloys Compod. 2010, 489, 269-

274.

[103] Farghali, A.A.; Moussa, M.; Khedr, M.H. J. Alloys Compd. 2010, 499, 98-103.

[104] Burianova, S.; Poltierova-Vejpravova, J.; Holec, P.; Plocek, J. J. Phys. Conf. Ser.

2010, 200, 072015.

[105] Reddy, M.P.; Madhuri, W.; Reddy, N.R.; Kumar, K.V.S.; Murthy, V.R.K.;

Reddy, R.R. Mater. Sci. Eng., C 2010, 30, 1094-1099.

[106] Kumar, G.; Chand, J.; Dogra, A.; Kotnala, R.K.; Singh, M. J. Phys. Chem. Solids

2010, 71, 375-380.

[107] Ghatak, S.; Sinha, M.; Meikap, A.K.; Pradhan, S.K. Mater. Res. Bull. 2010, 45,

954-960.

[108] Gabal, M.A.; Bayoumy, W.A. Polyhedron, 2010, 29, 2569-2573.

[109] Bharti, D.C.; Mukherjee, K.; Majumder, S.B. Mater. Chem. Phys. 2010, 120,

509-517.

[110] Ghasemi, A.; Ashrafizadeh, A.; Paesano Jr. A.; Machado, C.F.C.; Shirsath, S.E.;

Liu, X.; Morisako, A. J. Magn. Magn. Mater. 2010, 322, 3064-3071.

[111] Dalt, S.D.; Takimi, A.S.; Volkmer, T.M.; Sousa, V.C.; Bergmann, C.P.

Powder Technol. 2011, 210, 103-108.

[112] Bachhav, S.G.; Patil, R.S.; Ahirrao, P.B.; Patil, A.M.; Patil, D.R. Mater. Chem.

Phys. 2011, 129, 1104-1109.

[113] Sadhana, K.; Praveena, K.; Murthy, S.R. J. Magn. Magn. Mater. 2011, 323,

2977-2981.

157

[114] Mansour, S.F. J. Magn. Magn. Mater. 2011, 323, 1735-1740.

[115] Mansour, S.F.; Elkestawy, M.A. Ceram. Int. 2011, 37, 1175-1180.

[116] Khan, M.A.; Islam, M.U.; Ishaque, M.; Rahman, I.Z. Ceram. Int. 2011, 37, 2519-

2526.

[117] Kaiser, M. Physica B 2011, 406, 899-905.

[118] Shah, J.; Arora, M.; Purohit, L.P.; Kotnala, R.K. Sens. Actuators A 2011, 167,

332-337.

[119] Ferreira da Silva, M.G.; Valente, M.A. Mater. Chem. Phys. 2012, 132, 264-272.

[120] Patil, J.Y.; Khandekar, M.S.; Mulla, I.S.; Suryavanshi, S.S. Curr. Appl. Phys.

2012, 12, 319-324.

[121] Mounkachi, O.; Hamedoun, M.; Belaiche, M.; Benyoussef, A.; Masrour, R.; El

Moussaoui, H.; Sajieddine, M. Physica B 2012, 407, 27-32.

[122] Mohammed, K.A.; Al-Rawas, A.D.; Gismelseed, A.M.; Sellai, A.; Widatallah,

H.M.; Yousif, A.; Elzain, M.E.; Shongwe, M. Physica B 2012, 407, 795-804.

[123] Kaur, M.; Rana, S.; Tarsikka, P.S. Ceram. Int. 2012, 38, 4319-4323.

[124] Khan, M.A.; Islam, M.U.; Ishaque, M.; Rahman, I.Z. J. Alloys Compd. 2012,

519, 156-160.

[125] Singh, N.; Agarwal, A.; Sanghi, S.; Khasa, S. J. Magn. Magn. Mater. 2012, 324,

2506-2511.

[126] Sujatha, C.; Reddy, K.V.; Babu, K.S.; Reddy, A.R.C.; Rao, K.H. Physica B

2012, 407, 1232-1237.

[127] Albuquerque, A.S.; Tolentino, M.V.C.; Ardisson, J.D.; Moura, F.C.C.;

Mendonca, R.; Macedo, W.A.A. Ceram. Int. 2012, 38, 2225-2231.

[128] Chen, D.; Zhang, Y.; Tu, C. Mater. Lett. 2012, 82, 10-12.

[129] Bobade, D.H.; Rathod, S.M.; Mane, M.L. Physica B 2012, 407, 3700-3704.

[130] Sujatha, C.; Reddy, K.V.; Babu, K.S.; Reddy, A.R.C.; Rao, K.H. Ceram. Int.

2012, 38, 5813-5820.

158

[131] Vestal, C.R.; Zhang, J.Z. Int. J. Nanotechnol. 2004, 1, 240-263.

[132] Seip, C.T.; Carpenter, E.E.; O’Connor, C.J.; John, V.T.; Li, S. IEEE Trans.

Magn. 1998, 34, 1111-1113.

[133] Pillai, V.; Shah, D.O. J. Magn. Magn. Mater. 1996, 163, 243-248.

[134] Subbarao, M.V.; Yusuf, S.M.; Kulkarni, R.G.; Rao, L.M. Pramana-J. Phys. 1999,

53, 341-351.

[135] Shepherd, P.; Mallick, K.K.; Green, R.J. J. Magn. Magn. Mater. 2007, 311, 683-

692.

[136] Hamada, S.; Matijević, E. J. Colloid Interface Sci. 1981, 84, 274-277.

[137] Sugimoto, T.; Sakata, K.; Muramatsu, A. J. Colloid Interface Sci. 1993, 159, 372-

382.

[138] Burda, C.; Chen, X.; Narayana, R.; El-Sayed, M.A. Chem. Rev. 2005, 105, 1025-

1102.

[139] Rezlescu, N.; Doroftei, C.; Rezlescu, E.; Popa, P.D. J. Alloys Compd. 2008, 451,

492-496.

[140] Hua, Z.H.; Li, S.Z.; Han, Z.D.; Wang, D.H.; Lu, M.; Zhong, W.; Gu, B.X.; Du,

Y.W. Mater. Sci. Eng. A, 2007, 448, 326-329.

[141] Han, Z.; Dong, L.; Wu, Z.; Zhang, X. J. Rare Earths, 2006, 24, 75-77.

[142] Dong, L.; Han, Z.; Zhang, Y.; Wu, Z.; Zhang, X. Rare Met. 2006, 25, 605.

[143] Thompson, S.; Shirtcliffe, N.J.; O’Keefe, E.S.; Appleton, S.; Perry, C.C. J.

Magn. Magn. Mater. 2005, 292, 100-107.

[144] Haijun, Z.; Zhichao, L.; Chenliang, M.; Xi, Y.; Liangying, Z.; Mingzhong, W.

Mater. Chem. Phys. 2003, 80, 129-134.

[145] Ghasemi, A.; Morisako, A. J. Magn. Magn. Mater. 2008, 320, 1167-1172.

159

[146] Jotania, R.B.; Khomane, R.B.; Chauhan, C.C.; Menon, S.K.; Kulkarni, B.D. J.

Magn. Magn. Mater. 2008, 320, 1095-1101.

[147] Ataie, A.; Harris, I.R.; Ponton, C.B. in: J. David (Ed.), Introduction to

Magnetism and Magnetic Materials, Chapman & Hall, London, 1991, pp. 273.

[148] Pileni, M.P. J. Phys. Chem. 1993, 97, 6961-6973.

[149] Kulkarni, S.; Shrotri, J.; Deshpande, C.E.; Kate, S.K. J. Mater. Sci. 1989, 24,

3739-3744.

[150] Iida, H.; Takayanagi, K.; Nakanishi, T.; Osaka, T. J. Colloid Interface Sci. 2007,

314, 274-280.

[151] Hahn, H. Nanostruct. Mater. 1997, 9, 3-12.

[152] Byrappa, K.; Adschiri, T. Prog. Cryst. Growth Charact. Mater. 2007, 53, 117-

166.

[153] Kandori, K.; Ishikawa, T. J. Colloid Interface Sci. 2004, 272, 246-248.

[154] Konishi, Y.; Kawamura, T.; Asai, S. Metall. Mater. Trans. B, 1994, 25, 165-170.

[155] Courtecuisse, V.G.; Bocquet, J.F.; Chhor, K.; Pommier, C. J. Supercrit. Fluids,

1996, 9, 222-226.

[156] Bocquet, J.F.; Chhor, K.; Pommier, C. Mater. Chem. Phys. 1999, 57, 273-280.

[157] Thimmaiah, S.; Rajamathi, M.; Singh, N.; Bera, P.; Meldrum, F.;

Chandrasekhar, N.; Seshadri, R. J. Mater. Chem. 2001, 11, 3215-3221.

[158] Gautam, U.K.; Ghosh, M.; Rajamathi, M.; Seshadri, R. Pure Appl. Chem. 2002,

74, 1643-1649.

[159] Wang, J.F.; Ponton, C.B.; Grossinger, R.; Harris, I.R. J. Alloys Compd. 2004,

369, 170-177.

160

[160] Komarneni, S.; Menon, V.C.; Li, Q.H.; Roy, R.; Ainger, F. J. Amer. Ceram. Soc.

1996, 79, 1409-1412.

[161] Komarneni, S.; Arrigo, M.C.; Pellacani, G.C.; Katsuki, H. J. Amer. Ceram. Soc.

1998, 81, 3041-3043.

[162] Wang, Y.; Tang, X.; Yin, L.; Huang, W.; Gedanken, A. Adv. Mater. 2000, 12,

1137.

[163] Kodas, T.T. Adv. Mater. 1989, 6, 180-186.

[164] Senzaki, Y.; Caruso, J.; Hampden-Smith, M.J.; Kodas, T.T.; Wang, L.-M. J.

Am. Ceram. Soc. 1995, 78, 2973-2976.

[165] Shafi, K.V.P.M.; Koltypin, Y.; Gedanken, A.; Prozorov, R.; Balogh, J.;

Lendvai, J.; Felner, I. J. Phys. Chem. B 1997, 101, 6409-6414.

[166] Khalil, H.; Mahajan, D.; Rafailovich, M.; Gelfe, M.; Pandya, K. Langmuir, 2004,

20, 6896-6903.

[167] Lee, S.S.; Seo, K.W.; Yoon, S.H.; Shim, I.W.; Byun, K.T.; Kwak, H.-Y. Bull.

Korean Chem. Soc. 2005, 26, 1579-1581.

[168] Kim, H.W.; Kang, K.I.; Kwak, H.Y. Int. J. Hydrogen Energy, 2009, 34, 3351-

3359.

[169] Suslick, K.S.; Choe, S.B.; Cichowlas, A.A.; Grinstaff, M.W. Nature. 1991,

353, 414-416.

[170] Arean, C.O.; Blanco, J.L.R.; Gonzalez, J.M.R.; Fernandez, M.C.T. J. Mater.

Sci. Lett. 1990, 9, 229.

[171] Zhang, Z.J.; Wang, Z.L.; Chakoumakos, B.C.; Yin, J.S. J. Am. Chem. Soc.

1998, 120, 1800-1804.

[172] Turner, S.R.; Siano, D.B.; Bock, J.; U. S. Patent No. 4,521,580, 1985.

[173] Ovando, V.M. Polym. Bull. 2005, 54, 129-140.

161

[174] Boutonnet, M.; Kizling, J.; Stenius, P.; Marie, G. Colloids Surf. 1982, 5, 209-225.

[175] Pileni, M.P.; Fendler, J.H. Nanoparticles and Nanostructures Films, Wiley-

VCH, Weinheim, 1997.

[176] Li, Y.; Park, C.W. Langmuir, 1999, 15, 952-956.

[177] Iqbal, M.J.; Ahmad, Z.; Melikhov, Y.; Nlebedim, I.C. J. Magn. Magn. Mater.

2012, 324, 1088-1094.

[178] Hoar, T.P.; Schulman, J.H. Nature, 1943, 152, 102-103.

[179] Paul, B.K.; Moulik, S.P. Curr. Sci. 2001, 80, 990-1001.

[180] Paul, B.K.; Moulik, S.P. J. Dispersion Sci. Technol. 1997, 18, 301-367.

[181] Chen, J.P.; Lee, K.M.; Sorensen, C.M.; Klabunde, K.J.; Hadjipanayis, G.C. J.

Appl. Phys. 1994, 75, 5876-5878.

[182] Lisiecki, I.; Pileni, M.P. J. Am. Chem. Soc. 1993, 115, 3887-3896.

[183] Foos, E.E.; Stroud, R.M.; Berry, A.D.; Snow, A.W.; Armistead, J.P. J. Am.

Chem. Soc. 2000, 122, 7114-7115.

[184] Chew, C.H.; Gan, L.M.; Shah, D.O. J. Disp. Sci. Tech. 1990, 11, 593-609.

[185] Pang, Y.X.; Bao, X. J. Mater. Chem. 2002, 12, 3699-3704.

[186] Lu, C.-H.; Wang, H.-C. J. Mater. Chem. 2003, 13, 428-431.

[187] Xu, P.; Han, X.; Wang, M. J. Phys. Chem. 2007, 111, 5866-5870.

[188] Yue, Z.; Guo, W.; Zhou, J.; Gui, Z.; Li, L. J. Magn. Magn. Mater. 2004, 270,

216-233.

[189] Charsley, E.L.; Warrington, S.B. Thermal Analysis: Techniques and

Applications, Royal Society of Chemistry, Cambridge, 1992. Ch. 2.

162

[190] Giacovazzo, C. Fundamentals of Crystallography; 2nd

ed., Oxford University

Press, New York, 2002.

[191] Cullity, B.D. Elements of X-ray Diffraction, 3rd

ed., Addsion-Wesley Publishers,

New York, 1978, ch. 4.

[192] West, A.R. Basic Solid State Chemistry, 2nd ed., John Wiley & Sons, New York,

1999, pp. 126.

[193] Langford, J.I.; Lauër, D. Rep. Prog. Phys. 1996, 59, 131-234.

[194] Xu, G.; Ma, H.; Zhong, M.; Zhou, J.; Yue, Y.; He, Z. J. Magn. Magn. Mater.

2006, 301, 383-388.

[195] Feng, Q.; Jen, L. IEEE Trans. Magn. 2002, 38, 1391-1394.

[196] Zhou, W.; Wang, Z.L. Scanning Microscopy for Nanotechnology: Techniques

and Applications, Springer, New York, 2006.

[197] Whiston, C. X-ray Methods, John Wiley & Sons, New York, 1987, pp. 166-174,

and 297-301.

[198] Greenwood, N.N.; Gibb, T.C. Mössbauer Spectroscopy, Chapman and Hall,

London, 1971, p. 259.

[199] Wertheim, G.K. Mössbauer Effect: Principles and Applications, Academic Press,

New York, 1964, pp. 42.

[200] Bhide, V.G. Mössbauer Effect and its Applications, Tata McGraw Hill, New

Delhi, 1973.

[201] Clauser, M.J.; Mössbauer, R.L. Phys. Rev. 1969, 178, 559-567.

[202] Bancroft, G.M. Mössbauer Spectroscopy: An Introduction to Inorganic Chemists

and Geochemists, McGraw Hill, London, 1973, pp. 252.

[203] Goldanskii, V.I.; Herber, R.H. Chemical Applications of Mössbauer

Spectroscopy, Academic Press, New York, 1968, pp. 103.

163

[204] Longworth, G.; Window, B. J. Phys. D, 1971, 4, 835-839.

[205] Bancroft, G.M. Mössbauer Spectroscopy: An Introduction for Inorganic

Chemists and Geochemists, Wiley and Sons, New York, 1973, pp. 251.

[206] Margeat, O.; Amiens, C.; Chaudret, B.; Lecante, P.; Benfield, R.E. Chem. Mater.

2005, 17, 107-111.

[207] Zitoun, D.; Respaud, M.; Fromen, M.C.; Casanove, M.J.; Lecante, P.; Amiens,

C.; Chaudret, B. J. Phys. Chem. B 2003, 107, 6997-7005.

[208] Respaud, M.; Broto, J.M.; Rakoto, H.; Fert, A.R. Phys. Rev. B 1998, 57, 2925-

2935.

[209] Mahdi, A.E.; Panina, L.; Mapps, D. Sens. Actuators, A 2003, 105, 271-285.

[210] Gramm, K.; Lundgren, L.; Beckman, O. Phys. Scr. 1976, 13, 93-95.

[211] Chikazumi, S. Physics of Ferromagnetism, 2nd ed., Oxford University Press,

Oxford, 1997, pp. 503.

[212] Lee, E.W. Rep. Prog. Phys. 1955, 18, 184-229.

[213] Bozorth, R.M. Ferromagnetism, IEEE Press, New York, 1993, pp. 486.

[214] Dodrill, B.C. Low Moment Measurements with a Vibrating Sample

Magnetometer, LakeShore Cryotronics, Inc. 575 McCorkle Blvd. Westerville,

Ohio, 2004.

[215] VSM System-Hardware Reference Manual for 7400 Series, LakeShore

Cryotronics, Inc. 575 McCorkle Boulevard Westerville, 2004.

[216] Kojima, H. Ferromagnetic Materials: A Handbook on the Properties of

Magnetically Ordered Substances, North-Holland, Amsterdam, 1982, pp. 318.

[217] Iqbal, M.J.; Farooq, S. Mater. Res. Bull. 2009, 44, 2050-2055.

[218] Kao, K.C. Dielectric Phenomena in Solids, Elsevier Academic Press, California,

2004, pp. 51.

164

[219] Sun, T.; Borrasso, A.; Liu, B.; Dravid, V. J. Am. Ceram. Soc. 2011, 94, 1490-

1495.

[220] Zhu, J.; Gui, Z.; Ding, Y. Mater. Lett. 2008, 62, 2373-2376.

[221] Iqbal, M.J.; Ahmad, Z. J. Power Sources, 2008, 179, 763-769.

[222] Mishra, S.; Kundu, T.K.; Barick, K.C.; Bahadur, D.; Chakravorty, D. J. Magn.

Magn. Mater. 2006, 307, 222-226.

[223] Che, S.; Wang, J.; Chen, Q. J. Phys.: Condens. Matter. 2003, 15, L335-L339.

[224] Iqbal, M.J.; Ahmad, Z.; Meydan, T.; Melikhov, Y. J. Magn. Magn. Mater. 2012,

http://dx.doi.org/10.1016/j.jmmm.2012.06.031

[225] Lakshman, A.; Rao, P.S.V.S.; Rao, K.H. J. Magn. Magn. Mater. 2004, 284,

352-357.

[226] Soibam, I.; Phanjoubam, S.; Prakash, C. J. Alloys Compd. 2009, 475, 328-331.

[227] Wang, L.; Zhou, Q.; Li, F. Phys. Stat. Sol. (B) 2004, 241, 377-382.

[228] Roumaih, K.; Manapov, R.A.; Sadykov, E.K.; Pyataev, A.V. J. Magn. Magn.

Mater. 2005, 288, 267-275.

[229] Ashiq, M.N.; Naz, F.; Malana, M.A.; Gohar, R.S.; Ahmad, Z. Mater. Res. Bull.

2012, 47, 683-686.

[230] Pailhe, N.; Wattiaux, A.; Gaudon, M.; Demourgues, A. J. Solid State Chem. 2008,

181, 1040-1047.

[231] Isfahani, M.J.N.; Myndyk, M.; Šepelák, V.; Amighian, J. J. Alloys Compd. 2009,

470, 434-437.

[232] Thummer, K.P.; Chhantbar, M.C.; Modi, K.B.; Baldha, G.J.; Joshi, H.H. J.

Magn. Magn. Mater. 2004, 280, 23-30.

165

[233] Krieble, K.; Lo, C.C.H.; Melikhov, Y.; Snyder, J.E. J. Appl. Phys. 2006, 99,

08M912.

[234] Lee, S.J.; Lo, C.C.H.; Matlage, P.N.; Song, S.H.; Melikhov, Y.; Snyder, J.E.;

Jiles, D.C. J. Appl. Phys. 2007, 102, 073910.

[235] Cullity, B.D. Introduction to Magnetic Materials, Reading, MA: Addison

Wesley, 1972, p. 233.

[236] Ranvah, N.; Melikhov, Y.; Jiles, D.C.; Snyder, J.E.; Moses, A.J.; Williams, P. I.;

Song, S.H. J. Appl. Phys. 2008, 103, 07E506.

[237] Shenker, H. Phys. Rev. 1957, 107, 1246-1249.

[238] Slonczewski, J.C. Phys. Rev. 1958, 110, 1341-1348.

[239] Sakurai, S.; Sasaki, S.; Okube, M.; Ohara, H.; Toyoda, T. Physica B, 2008, 403,

3589-3595.

[240] Zaki, H.M. Physica B, 2009, 404, 3356-3362.

[241] El Guendouzi, M.; Sbai, K.; Perriat, P.; Gillot, B. Mater. Chem. Phys. 1990, 25,

429-436.

[242] Ranvah, N.; Nlebedim, I.C.; Melikhov, Y.; Snyder, J.E.; Williams, P.I.; Moses,

A.J.; Jiles, D.C. IEEE T. Magn. 2009, 45, 4261-4264.

[243] Tachiki, M. Prog. Theor. Phys. 1960, 23, 1055-1072.

[244] Ranvah, N.; Melikhov, Y.; Nlebedim, I.C.; Jiles, D.C.; Snyder, J.E.; Moses, A.J.;

Williams, P.I. J. Appl. Phys. 2009, 105, 07A518.

[245] Ghasemi, A.; Morisako, A. J. Alloys Compd. 2008, 456, 485-491.

[246] Ghasemi, A.; Hossienpour, A.; Morisako, A.; Saatchi, A; Salehi, M. J. Magn.

Magn. Mater. 2006, 302, 429-435.

[247] Li, Y.; Liu, R.; Zhang, Z.; Xiong, C. Mater. Chem. Phys. 2000, 64, 256-259.

166

[248] Goldman, A. Modern Ferrite Technology, II, Springer Science Inc., USA, 2006,

pp. 14.

[249] Van der Zaag, P.J.; Brabers, V.A.M.; Johnson, M.T.; Noordemeer, A.; Bongers,

P.F. Phys. Rev. B, 1995, 51, 12009-12011.

[250] Tang, Z.X.; Sorensen, C.M.; Klabunde, K.L.; Hadjipanayis, G.C. Phys. Rev. Lett.

1991, 67, 3602-3605.

[251] Melagiriyappa, E.; Jayanna, H.S. J. Alloys Compd. 2009, 482, 147-150.

[252] Chhaya, U.V.; Kulkarni, R.G. Mater. Lett. 1999, 39, 91-96.

[253] Rezlescu, E.; Rezlescu, N.; Popa, P.D. Phys. Stat. Sol. 2004, 1, 3636-3639.

[254] Zahi, S.; Daud A.R.; Hashim, M. Mater. Chem. Phys. 2007, 106, 452-456.

[255] Ellis, A.B. J. Chem. Education, 1987, 64, 836-846.

[256] Ponpandian, N.; Balaya, P.; Narayanasamy, A. J. Phys.: Condens. Mater. 2002,

14, 3221-3237.

[257] Roy, P.K.; Bera, J. J. Mater. Process. Tech. 2008, 197, 279-283.

[258] Iqbal, M.J.; Siddiquah, M.R. J. Magn. Magn. Mater. 2008, 320, 845-850.

[259] Gul, I.H.; Amin, F.; Abbasi, A.Z.; Rehman, M.A.; Maqsood, A. Scripta Mater.

2007, 56, 497-500.

[260] Ashiq, M.N.; Saleem, S.; Malana, M.A.; Rehman, A. J. Alloys Compd. 2009, 486,

640-644.

[261] Pandya, M.P.; Modi, K.B.; Joshi, H.H. J. Mater. Sci. 2005, 40, 5223-5232.

[262] Ahmed, M.A.; Ateia, E.; El-Dek, S.I. Mater. Lett. 2003, 57, 4256-4266.

[263] Rana, C.P.; Baijal, J.S.; Kishan, P. J. Less-Common Met. 1983, 165, 257-261.

167

[264] Zhang, H.; Zhou, J.; Wang, Y.; Li, L.; Yue, Z.; Gui, Z. Mater. Lett. 2002, 55, 351-

355.

[265] Iqbal, M.J.; Farooq, S. Mater. Chem. Phys. 2009, 118, 308-313.

[266] Ahmed, M.A.; Okasha, N.; El-Sayed, M.M. Ceram. Int. 2007, 33, 49-58.

[267] Kaiser, M. J. Alloys Compd. 2009, 468, 15-21.

[268] Shinde, T.J.; Gadkari, A.B.; Vasambekar, P.N. Mater. Chem. Phys. 2008, 111,

87-91.

[269] Iqbal, M.J.; Yaqub, N.; Sepiol, B.; Ismail, B. Mater. Res. Bull. 2011, 46, 1837-

1842.

[270] Verma, A.; Goel, T. C.; Mendiratta, R. G.; Gupta, R. G. J. Magn. Magn. Mater.

1999, 192, 271-276.

[271] Wang, Y.Z.; Qiao, G.W.; Liu, X.D.; Ding, B.Z.; Hu, Z.Q. Mater. Lett. 1993, 17,

152-154.

[272] Liu, X.D.; Ding, B.Z.; Hu, Z.Q.; Lu, K.; Wang, Y.Z. Physica B, 1993, 192,

345-350.

[273] Lee, H.Y.; Burton, L.C. IEEE Trans. Comp. Hybrids Manuf. Tech. 1987, 482,

712-718.

[274] Wang, Z.H.; Ji, T.H.; Wang, Y.Q.; Chen, X.; Li, R.W.; Cai, J.W. J. Appl. Phys.

2000, 87, 5582.

[275] Iwauchi, K. Jpn. J. Appl. Phys. 1971, 10, 1520.

[276] Rezlescu, N.; Rezlescu, E. Phys. Status Solidi A, 1974, 23, 575-582.

[277] Koops, C.G. Phys. Rev. 1951, 83, 121-127.

[278] Narang, S.B.; Hudiara, I.S. J. Ceram. Proces. Res. 2006, 7, 113-116.

[279] Mallick, K.K.; Shepherd, P.; Green, R.J. J. Eur. Ceram. Soc. 2007, 27, 2045-

2052.

168

[280] El-Hiti, M.A. J. Magn. Magn. Mater. 1999, 192, 305-313.

[281] Iwauchi, K. Phys. Status Solidi A, 1985, 92, 309-314.

[282] Ashiq, M.N.; Iqbal, M.J.; Gul, I.H. J. Alloys Compd. 2009, 487, 341-345.

List of Publications

Iqbal, M.J.; Ahmad, Z.; Meydan, T.; Melikhov, Y. “Physical, electrical and

magnetic properties of nano-sized Co-Cr substituted magnesium ferrites” J. App.

Phys. 2012, 111, 033906.

Iqbal, M.J.; Ahmad, Z.; Meydan, T.; Nlebedim, I.C. “Influence of Ni-Cr

substitution on the magnetic and electric properties of magnesium ferrite

nanomaterials” Mater. Res. Bull. 2012, 47, 344-351.

Iqbal, M.J.; Ahmad, Z.; Melikhov, Y.; Nlebedim, I.C. “Effect of Cu-Cr co-

substitution on magnetic properties of nanocrystalline magnesium ferrite” J.

Magn. Magn. Mater. 2012, 324, 1088-1094.

Iqbal, M.J.; Ahmad, Z.; Meydan, T.; Melikhov, Y. “Temperature and

composition dependence of magnetic properties of cobalt-chromium co-

substituted magnesium ferrite” J. Magn. Magn. Mater. 2012, online.

The following contributions are not included in this dissertation:

Iqbal, M.J.; Ahmad, Z. “Electrical and dielectric properties of lithium manganate

nanomaterials doped with rare-earth elements” J. Power Sources, 2008, 179, 763-

769.

Ashiq; M.N.; Naz, F.; Malana, M.A.; Gohar, R.S.; Ahmad, Z. “Role of Co-Cr

substitution on the structural, electrical and magnetic properties of nickel

nanoferrites synthesized by the chemical co-precipitation method” Mater. Res.

Bull. 2012, 47, 683-686.

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