Fabrication, Characterization and Structural Study of

Fabrication, Characterization and Structural Study of Ferrites of Technical Importance

Ph.D

Shahid Mahmood Ramay

M.Phil. Session (2003-2008)

CENTRE OF EXCELLENCE IN SOLID STATE PHYSICS UNIVERSITY OF THE PUNJAB

LAHORE (PAKISTAN)

Fabrication, Characterization and Structural Study of Ferrites of Technical Importance

A thesis submitted, in partial fulfillment

of the requirement for the award of the degree of

DOCTOR OF PHILOSOPHY

In

SOLID STATE PHYISCS

By


M.Phil. Session (2003-2008)

Centre of Excellence in Solid State Physics University of the Punjab Quaid-i-Azam Campus

Lahore, Pakistan.

CERTIFICATE

This is to certify that research work contained in this thesis has been carried out by

Shahid Mahmood Ramay S/o Ch. Abdul Khaliq, Session (2003-2008), as partial

requirement for the award of degree of Ph.D (Solid State Physics). He is allowed to

submit this thesis to Centre of Excellence in Solid State Physics, University of Punjab,

Lahore.

Research Supervisor Director Dr. Saadat. A. Siddiqi, Dr. Shahzad Naseem, Professor, Professor and Director, Centre of Excellence in Centre of Excellence in Solid State Physics Solid State Physics University of the Punjab, University of the Punjab, Lahore. Lahore.

Acknowledgement In the name of Allah, the most gracious and merciful who bestowed me with

wisdom. I pay my humble salutations to my beloved The Holy Prophet,

Muhammad (Peace Be Upon Him).

I would like to express my sincerest thanks to my supervisor Prof. Dr. S.A.

Siddiqi for his guidance, support (moral and financial), cooperation and sympathetic

attitude during my PhD journey. I have got a profitable knowledge from his insight

and research experience. The completion of this thesis was not possible without his

guidance and support.

I am thankful to Director, Prof. Dr. Shahzad Naseem for his kind help and

suggestions. I am also thankful to Dr. Saira Riaz who helped me in XRD, VSM and

SEM labs. Thanks must be given to Prof. Dr. Sabieh Anwar (LUMS) for useful

suggestions and discussions.

I am thankful to HEC (Higher Education Commission) who provided me

fellowship of six months for South Korea. I extend my thanks to Prof. Dr. Sung

Chull Shin (Korean Advanced Institute of Science and Technology) who provides me

good opportunities in his laboratory. I gained a lot of knowledge and experience under

his kind supervision during my stay at South Korea.

Credit must be given to my lab. fellows, Shahid Atiq, Furrukh Shahzad,

Murtaza Saleem, M. Tanveer, Zafar Bhai and Huma Asif for their nice company,

suggestions and assistance. Outside the department, I would like to thank Prof. Dr.

Nizami (COMSAT Sahiwal), Anwar-ul-Haq (University of Lahore) and M. Ansar

(PCSIR Lab.) for their kind help.

I am also thankful to my elder brothers Tariq Mahmood, Dr. Zahid

Mahmood and brothers in law Tariq Majeed, Tahir Majeed and Talha Majeed

who helped me and prayed for me.

Finally, I want to thank my wife, who inspired me very much. Without her

patience, I would have never accomplished this task.


Dedicated

To

My Parents (Late)

&

My Wife

Abstract Ferrites are widely used in power electronics applications where the frequency

range is from KHz to MHz. No other alternative materials except ferrites are available

at such high frequencies. The areas of magnetic nanoparticles and thin films lead to

revolutionary new approaches in basic and advanced magnetism, and are more

effective in the field of high density storage media. The main objective of the present

study was to produce single phase ferrites in the form of bulk, nano and thin films

with improved structural, electrical and magnetic properties.

This thesis examines the issue encountered in the growth, structural,

microstructural, electrical and magnetic properties of ferrites in the form of bulk,

nanoparticles and thin films. Here the materials examined include Cu0.5Zn0.5Fe2-

xAlxO4 (x=0.0 to 0.5) ferrites prepared with solid state reaction method,

Co0.5Mn0.5Fe2O4 (calcined at 500, 600, 700, 800, 900°C), Mn0.5Cu0.5-xZnxFe2O4

(x=0.0 to 0.5), Mn0.5Cu0.5-xNixFe2O4 (x=0.0 to 0.5) ferrites prepared with sol-gel

combustion method and Fe3O4

The effect of Al

thin films prepared with pulsed laser deposition

technique. 3+ on the structural, electrical and magnetic properties were

investigated in Cu0.5Zn0.5Fe2-xAlxO4 (x=0.0 to 0.5) ferrites prepared with solid state

reaction method. Single phase cubic spinel structure was revealed by X-ray diffraction

analysis. For all the samples, crystallite size remained in the range of 25-30 nm.

Lattice constants of all the samples decreased, whereas porosity increased with

increasing Al+3 concentration due to the substitution of smaller Al3+ ion (0.51 Å) for

large Fe3+ ion (0.64 Å). Due to non-magnetic trend of Al3+ concentrations for a

magnetic element Fe3+ at the B-site gradually decreased the saturation magnetization.

Al+3 ε has significant impact on the dielectric constant ( /

ε

), tangent of dielectric loss

angle (tanδ) and dielectric loss factor ( //

Three series of ferrites Co

). The possible reason for the variation in

dielectric properties has been understood on the basis of space charge polarization.

0.5Mn0.5Fe2O4 (calcined at 500, 600, 700, 800,

900°C), Mn0.5Cu0.5-xZnxFe2O4 (x=0.0 to 0.5), Mn0.5Cu0.5-xNixFe2O4 (x=0.0 to 0.5)

were prepared by sol-gel combustion method. In Co0.5Mn0.5Fe2O4 ferrites, crystallite

size was determined with Scherrer’s formula. Crystallite size increases with

calcination temperature but coercivity decreases. The decrease in coercivity at larger

crystallite size can be attributed to domain walls. Single phase nanocrystalline

Mn0.5Cu0.5-xZnxFe2O4 (x=0.0 to 0.5) ferrites were successfully prepared at low

temperature of 300°C using citric acid as a fuel and nitrates as oxidants by sol-gel

method. X-ray diffraction (XRD) and room temperature vibrating sample

magnetometer (VSM) studies have been carried out in order to understand the

structural and magnetic properties as a function of zinc concentration. The variations

of observed lattice parameter and crystallite size have been explained by considering

the larger ionic radius of zinc. The coercivity decreases as the crystallite size increases,

attaining a minimum value of 46.32 Oe. This decrease at larger crystallite size could

be due to three reasons. First, the crossover of single domain to multiphase domain,

second combined effect of surface and surface anisotropy, third migration of Fe+3 ions

from A to B-site. Another series of single phase nano-crystalline Mn0.5Cu0.5-

xZnxFe2O4 (x=0.0 to 0.5) ferrites were successfully synthesized by combustion

method at a temperature as low as 300°C. The presence of Ni2+ ions did not show a

consistent trend in diffraction peaks shifting to either lower or higher angles. It was

observed that with increasing nickel concentration, saturation magnetization (Ms)

increased but coercivity (Hc) decreased which could be attributed to the substitution

of soft ferromagnetic Ni2+ ions in place of diamagnetic Cu2+ ions. The minimum value

of coercivity (87.20 Oe) was observed for the composition Mn0.5Ni0.5Fe2O4

Fe

.

3O4 thin films were deposited on Si(100) substrates with pulsed laser

deposition technique. First we studied the effect of annealing and deposition

temperature, and second the effect of annealing time of 30, 60 and 90 minutes on the

structural and magnetic properties of Fe3O4 thin films. Scanning electron microscopy,

X-ray diffractometery and vibrating sample magnetometry were used to find the film

thickness, Fe3O4 phase and magnetic properties respectively. We demonstrate

optimized deposition and annealing condition for an enhanced magnetization of 854

emu/cc that is very high as compared to the bulk sample. Effect of annealing time on

Fe3O4 thin films were studied by X-ray diffractometer and vibrating sample

magnetometer. Single phase [111] oriented Fe3O4

thin films independent of substrate

orientation was obtained after ninety minutes annealing. This preferred [111] oriented

growth was explained on the basis of the achievement of a thermodynamic stable state.

Table of Contents

Chapter-1

1.1 Significance of present work ...................................................................... 1

Introduction

1.2 Brief history and origin of magnetism ....................................................... 2

1.3 Magnetic materials ..................................................................................... 2

1.4 Classification of magnetic materials .......................................................... 3

1.5 Types of ferrites with respect to their magnetic properties ........................ 6

1.6 Types of ferrites with respect to their Structures ...................................... 7

1.6.1 Spinel cubic ferrites.................................................................................... 7

1.6.2 Spinel structure........................................................................................... 8

1.7 Types of spinel ferrites ............................................................................... 9

1.7.1 Substitutional ferrites ................................................................................. 10

1.8 Phases of Fe oxides .................................................................................... 10

1.9 Interactions in ferrimagnetics ..................................................................... 13

1.10 Thin film ferrites ........................................................................................ 14

1.11 Magnetic nanoparticles .............................................................................. 15

1.12 Applications of ferrites ............................................................................... 15

1.13 Densities of ferrites .................................................................................... 16

1.14 Porosity in ferrites ...................................................................................... 16

1.15 Hardness of ferrites .................................................................................... 16

1.16 Dielectric behavior of ferrites .................................................................... 17

1.17 Electrical resistance of ferrites ................................................................... 17

1.18 Magnetic behavior of ferrites ..................................................................... 17

References ................................................................................................. 19

Chapter-2

2.1 Cu-Zn ferrite .............................................................................................. 22

Literature Survey

2.2 Co-Mn ferrites ........................................................................................... 24

2.3 Mn-Cu ferrites ........................................................................................... 26

2.4 Fe3O4

References .................................................................................................................. 30

thin films ........................................................................................ 27

Chapter-3

3.1 Preparation Methods .................................................................................. 33

Experimental Techniques

3.1.1 Solid state reaction method .................................................................... 34

3.1.2 Sol-gel combustion method ................................................................... 35

3.1.3 Pulsed laser deposition technique .......................................................... 36

3.2 Characterization techniques ................................................................... 38

3.2.1 X-ray diffraction (XRD) ....................................................................... 38

3.2.2 Powder Method ..................................................................................... 39

3.2.3 Measurement of bulk density ................................................................. 40

3.2.4 Scanning electron microscopy (SEM) .................................................. 41

3.2.5 Vibrating sample magnetometer (VSM) ............................................... 42

3.2.6 Dielectric properties measurement ........................................................ 43

3.2.7 Electrical properties measurement ......................................................... 44

References ............................................................................................. 46

Chapter-4 Structural, magnetic and electrical properties of Al3+

4.1 Motivation ............................................................................................. 48

substituted Cu-Zn-ferrites

4.2 Sample preparation ................................................................................ 49

4.3 Results and discussion ........................................................................... 50

4.4 Conclusions ........................................................................................... 57

References ............................................................................................ 58

Chapter-5

5.1 Influence of temperature on the structural and magnetic properties of

Co

Fabrication and characterization of nanostructured

magnetic materials

0.5Mn0.5Fe2O4

5.1.1 Motivation ............................................................................................. 59

ferrites ............................................................ 59

5.1.2 Experimental details ............................................................................. 60

5.1.3 Results and discussions ......................................................................... 60

5.1.4 Conclusions ........................................................................................... 64

References ............................................................................................ 66

5.2 Low temperature synthesis of nanocrystalline Mn-Cu-Zn ferrties via sol-

gel combustion method .......................................................................... 68

5.2.1 Motivation .............................................................................................. 68

5.2.2 Experimental details............................................................................... 69

5.2.3 Results and discussion ........................................................................... 70

5.2.4 Conclusions ............................................................................................ 73

References ............................................................................................. 74

5.3 Low temperature synthesis and magnetic properties of Mn0.5Cu0.5-x

NixFe2O4

5.3.1 Motivation .............................................................................................. 75

nanoparticles via sol-gel combustion method ....................... 75

5.3.2 Experimental .......................................................................................... 76

5.3.3 Results and discussion ........................................................................... 77

5.3.4 Conclusions ............................................................................................ 80

References ............................................................................................. 82

Chapter-6 Fe3O4

6.1 Effect of temperature on structural and magnetic properties of laser

ablated iron oxide deposited on Si(100) ............................................... 83

thin films on Si(100) substrate with pulsed laser deposition

technique

6.1.1 Motivation ............................................................................................. 83

6.1.2 Preparation of Fe3O4

6.1.3 Characterizations.................................................................................... 85

thin films ............................................................. 84

6.1.4 SEM for thin-film thickness determination ........................................... 85

6.1.5 X-ray diffraction (XRD) analysis .......................................................... 86

6.1.6 Magnetic properties ............................................................................... 89

6.1.7 Conclusions ............................................................................................ 93

References ........................................................................................... 100

6.2 Effect of annealing time on structural and magnetic properties of laser

ablated oriented Fe3O4

6.2.1 Motivation .......................................................................................... ...95

thin films deposited on Si(100)

6.2.2 Preparation .......................................................................................... ...95

6.2.3 Results and discussion ........................................................................ ...96

6.2.4 Conclusions ......................................................................................... ...98

References ......................................................................................... .100

Conclusions .......................................................................... .101

Appendix ........................................................................................... .104

Fig.No. Description Page No.

List of Figures

1.1 Flux density in diamagnetic sample (Lines represent magnetic field lines) .. 3

1.2 Flux density in paramagnetic sample (Lines represent magnetic field lines) 4

1.3 Ferromagnetism (arrows represent atomic magnetic dipoles) ....................... 4

1.4 Antiferromagnetism (arrows represent atomic magnetic dipoles) ................. 4

1.5 Antiferromagnetic material with and without applied magnetic field (arrows

represent atomic magnetic dipoles) ......................................................... ….5

1.6 Ferrimagnetism (arrows represent atomic magnetic dipoles) ................... ….5

1.7 Ferrimagnetic material with and without applied magnetic field (arrows

represent atomic magnetic dipoles) .............................................................. 5

1.8 Plot of coercivity as a function of particle diameter ..................................... 6

1.9 Location of A and B-sites in unit cell ........................................................... 8

1.10 Local atomic arrangement for (a) tetrahedral site (b) octahedral site in spinel

structure......................................................................................................... 8

1.11 Crystal structure for different phases of iron oxides (a) Wustite (b) Hematite

(c) Magnetite and Maghemite ....................................................................... 11

1.12 Schematic representation of density of states (DOS) for a) Normal non

magnetic metal b) Normal ferromagnetic metal c) Half- metallic ferromagnetic

materials ........................................................................................................ 12

1.13 Crystallographic and magnetic structure in Fe3O4

3.1 Pestles and mortars for fine grinding ............................................................ 34

, near tetrahedral and

octahedrally (site A) and octahedrally (site B) coordinated Fe atoms. Here SE

represents “Superexchange” and DE represents “Double Exchange” ......... 14

3.2 Selection of porcelain, alumina or platinum crucibles.................................. 35

3.3 Combustion process ...................................................................................... 35

3.4 Thin film deposition under PLD-system....................................................... 36

3.5 Geometrical description of Bragg’s law ....................................................... 39

3.6 Schematic diagram of scanning electron microscopy ................................... 41

3.7 Schematic diagram of vibrating sample magnetometer ................................ 43

4.1 XRD patterns of Cu0.5Zn0.5AlxFe2-xO4

4.2 SEM micrographs of Cu

ferrite samples (x=0.0 to 0.5)......... 50

0.5Zn0.5AlxFe2-xO4

(b) x = 0.1, (c) x =0.2, (d) x = 0.3, (e) x = 0.4 and (f) x = 0.5 ..................... 52

with (a) x = 0.0,

4.3 Saturation magnetization plotted against Al3+

4.4 DC electrical resistivity plotted against temperature .................................... 54

concentration ....................... 53

4.5 Dielectric constant plotted against frequency ............................................... 55

4.6 Tangent of dielectric loss angle plotted against frequency. ......................... 56

4.7 Dielectric loss factor plotted against frequency ............................................ 56

5.1 XRD patterns of Co0.5Mn0.5Fe2O4

5.2 Variation of lattice parameters with calcination temperatures ..................... 62

, as-burnt and calcined at 500,600, 700, 800

and 900°C ..................................................................................................... 61

5.3 Variation of crystallite size with calcination temperature ............................ 62

5.4 Room temperature magnetic properties of Co0.5Mn0.5Fe2O4

5.5 Variation of coercivity with calcination temperature. ................................. 64

calcined at

different temperatures ................................................................................... 63

5.6 Coercivity as a function of crystallite size .................................................... 64

5.7 XRD patterns of Mn0.5Cu0.5-xZnxFe2O4

5.8 Variation of crystallite size and lattice parameter with zinc concentration of

Mn

ferrites ........................................... 70

0.5Cu0.5-xZnxFe2O4

5.9 Room temperature hysteresis loops of Mn

ferrites ...................................................................... 71

0.5Cu0.5-xZnxFe2O4

5.10 Variation of saturation magnetization (M

ferrites ........ 71

s) and coercivity (Hc

5.11 X-ray diffraction patterns of as-burnt Mn

) as a function of

zinc concentration ......................................................................................... 72

0.5Cu0.5-xNixFe2O4

5.12 Variation of lattice constant and crystallite size with Ni concentration ....... 78

powders ........ 77

5.13 RT hysteresis loops for Mn0.5Cu0.5-xNixFe2O4

5.14 Variation of coercivity and saturation magnetization as function of Ni

concentration ................................................................................................. 79

ferrites with varying Ni

concentration ................................................................................................. 79

5.15 Variation of coercivity and saturation magnetization as function of Ni

Concentration ................................................................................................ 80

6.1 SEM images of (a) Fe3O4

6.2 X-ray diffraction patterns of Fe

thin film annealed at 450°C and (b) as-deposited

film at 450°C ................................................................................................. 85

3O4 thin films on Si(100) substrates deposited

at room temperature and annealed at the shown temperatures ..................... 87

6.3 XRD diffraction patterns of as deposited thin films ..................................... 88

6.4 Inplane magnetization curves of annealed thin films with A, B and C

representing samples annealed at 300, 400 and 450°C ................................. 90

6.5 Inplane magnetization curves of as-deposited thin films with A, B and C

representing deposition temperatures of 350, 400 and 450°C ...................... 91

6.6 Variation of Hc

6.7 Variation of H

with crystallite size of annealed thin films.......................... 92

c

6.8 Film thickess measured by scanning electron microscopy (SEM) .............. 96

with crystallite size of as-deposited thin films .................... 92

6.9 XRD-patterns of Fe3O4

6.10 Vibrating sample magnetometry(VSM) of Fe

thin films at different annealing time .................... 97

3O4

thin films annealed at 0, 30,

60 and 90 minutes ......................................................................................... 98

List of Tables Table No Description

3.1 Compositions with their concentrations/calcinations temperature and

preparation techniques .................................................................................. 33

3.2 Pulsed Nd:YAG Laser NL303 (EKSPLA) Specifications............................ 37

3.3 Diffraction method with their wavelengths and angle .................................. 39

4.1 Lattice constant (a), lattice volume (V), sintered density (ρs), X-ray density

(ρx), porosity (P), saturation magnetization (Ms), activation energy ( E) of

Cu0.5Zn0.5Fe2-xAlxO2

5.1 Crystallite size, lattice constants, coercivity and magnetization of Mn

ferrite system ............................................................. 51

0.5Cu0.5-

xZnxFe2O4

5.2 Crystallite size, lattice constant, coercivity and magnetization of Mn

ferrites ........................................................................................ 73

0.5Cu0.5-

xNixFe2O4

6.1 Target materials with deposition and annealing temperature and time ........ 85

ferrites ......................................................................................... 80

6.2 XRD and VSM analysis of annealed Fe3O4

Si(100) substrates .......................................................................................... 89

thin films on

6.3 XRD and VSM analysis of as-deposited Fe3O4

on Si(100) substrates ..................................................................................... 89

thin films

6.4 Lattice parameter, crystallite size, lattice strain, saturation magnetization and

coercivity of all samples ............................................................................... 97

Chapter 1 Introduction

1

Introduction

The present work details an investigation into several materials systems with a

focus on their applicability for solid state electronic and microelectronic industry.

This work basically explains the growth and properties of the materials such as

chemistry, crystal structure, magnetization and electrical characterizations. Here

material systems examined include Fe3O4 thin films, Co0.5Mn0.5Fe3O4 nanoparticles,

Mn0.5Cu0.5-x ZnxFe2O4 nanoparticles, Mn0.5Cu0.5-x NixFe2O4 nanoparticles and

Cu0.5Zn0.5Fe2-x AlxO4 ferrites,

1.1 Significance of present work

prepared with pulsed laser deposition, sol-gel auto-

combustion and usual solid state reaction methods respectively.

The advancement in the electronics and microelectronic industries are due to

development of new materials. Advanced magnetic materials are one of them with

broad range of applications including data storage and circuit components such as

inductors and transformers. Today the general trend in electronics is toward magnetic

thin films and more powerful devices such as we can see in microprocessor speed and

magnetic data storage.

The reason for the selection of ferrites as a research topic has been due to their

immensely rich structural and magnetic properties and their vast technical and

industrial applications. Research on ferrites offers an excellent chance to explore

various aspects affecting their overall structural and magnetic properties. One of the

main emphases of the present work would be to study the relationship between

structural parameters and different concentrations of the substituted magnetic and non

magnetic ions prepared by pulsed laser deposition, sol-gel auto-combustion and

standard ceramic method. During the last few years, many countries have made great

progress in preparing magnetic oxides, which can be used in simple home appliances

to space technology.

Unfortunately, in Pakistan in spite of vast applications, we could not make progress in

the production of suitable ferrite magnetic materials. Sufficient technical know how

and basic science is lacking in the development of such an important materials. The

present study is an attempt to develop sufficient understanding for the production of

these materials by exploiting locally available indigenous sources.


2

1.2 Brief history and origin of magnetism “Magnetism” is a phenomenon in which some materials exert a force of

attraction or repulsion on other materials. This term came from Magnesia, an Island in

Aegean Sea, where certain stones were found by the Greeks in 470 BC. These stones

are “Lodestone” or magnetite (FeO.Fe2O3

1.3 Magnetic materials

) i.e, ferroferrite. Many of our modern

technological devices rely on magnetism and magnetic materials; these include

electrical and power generators and transformers, electric motors, radio, television,

telephones, computers and components of sound and video reproduction systems.

William Gilbert (1540-1603) was the first person who studied scientifically

magnetism and published his classic book on the Magnetism in 1600. He made some

experiments with loadstones and irons magnets, and gave a clear picture of the Earth’s

magnetic field. In eighteen century compound steel magnets were made, composed of

many magnetized steel strips fastened together, which could lift 28 times their own

weight of iron. In 1820, Oersted (1775-1851) performed an experiment in which he

produced magnetic field with electric current and this phenomena is called

Electromagnetism [1].

Magnetic materials are playing a crucial and major role in many devices of

every-day life: ac and dc motors, power distribution systems, based on power

transformers, which deliver energy for home and industrial use; video and audio

applications which provide information and entertainment on a massive scale;

telephone and telecommunication systems (microwave devices) which link continents

at nearly the speed of light; data storage systems (discs, disc drives) which pervade

virtually every human activity [2].

Soft magnetic materials can be easily magnetized and demagnetized. They are used

in applications such as cores of distribution, power transformers, small electric

transformers and stator, and rotor materials for motors and generators [3].

Permanent magnets are referred to as hard magnets. They are used as permanent

magnets in loud-speakers, telephone receivers, synchronous and brushless motors,

automotive starting motors [3].


3

1.4 Classification of magnetic materials

Magnetic properties originate from magnetic moment of the constituent atoms

which are produced by spin and orbital motion of their electrons. Most of the

atoms have completely paired electrons i.e, for each electron spinning in one

direction, there is another electron spinning in the opposite direction. The

same situation exists in orbital motion of the electrons. So the net circulating

current about any axis is zero exhibiting zero magnetic moment. These

substances show very weak magnetic behavior and are called non-magnetic.

However vacuum is only truly non-magnetic medium.

Experimentally and theoretically, all matter may be classified into following

groups;

1. Diamagnetic materials

2. Paramagnetic materials

3. Ferromagnetic materials

4. Antiferromagnetic materials

5. Ferrimagnetic materials

6. Superparamagnetic materials

Diamagnetic materials

Diamagnetic materials are those which, when placed in a magnetic field,

becomes weakly magnetized. The resultant magnetization is opposed to the applied

field. These materials contain no dipoles, but only dipoles that are induced by an

external field. Such materials show repulsion and negative susceptibility to external

magnetic field. Their susceptibility does not depend on temperature. Some examples

of diamagnetic materials are gold, copper, silver and superconductors [4]

Fig. 1.1 Flux density in diamagnetic sample (Lines represent magnetic field lines) [5]

Paramagnetic materials

Paramagnetic materials are those which, when placed in a magnetic field,

becomes weakly magnetized in the direction of the applied field. Paramagnetic

materials contain permanent dipoles. These materials show positive magnetization


4

and susceptibility. Also they are attracted by the magnetic field and the magnetic

properties are not retained after the field is removed. Susceptibility of these materials

depends on temperature which can be explained on Curie law. Examples of

paramagnetic materials are Li, Mo and Mn [4].

Fig. 1.2 Flux density in paramagnetic sample (Lines represent magnetic field

lines) [5]

Ferromagnetic materials

Ferromagnetic materials are those which, when placed in a magnetic field,

becomes strongly magnetized in the direction of the applied field. These materials

have strong magnetic properties due to the presence of magnetic domain. Here

domains have large number of atoms (1012-1015) with magnetic moments aligned

parallel to each other. These materials have very strong interactions due to the

electronic exchange forces. The ferromagnetism disappears if the temperature is

increased above a critical value called Curie temperature and the substance becomes

paramagnetic. The Curie temperature is an intrinsic property of the material and we

can use this property to identify any magnetic material [4]

Fig. 1.3 Ferromagnetism (arrows represent atomic magnetic dipoles) [5]

Antiferromagnetic materials

Antiferromagnetic materials are those materials in which magnetic moments

are arranged into groups, which contribute equal and opposite net magnetization.

These materials can not have any magnetization in the absence of an applied field.

The opposite atomic magnetic moments are due to the quantum mechanical exchange

forces. Antiferromagnetism occurs below a certain temperature called ‘Neel

temperature’ (TN). Above this temperature, materials become paramagnetic [6].


5

Fig. 1.4 Antiferromagnetism (arrows represent atomic magnetic dipoles) [5]

Fig. 1.5 Antiferromagnetic material with and without applied field

(arrows represent atomic magnetic dipoles)

Ferrimagnetic materials

Ferrimagnetic materials as shown in Fig.1.7 are called ferrites and are

technically important for data storage devices and power electronics in the form of

thin films, nano particles and bulk materials. Ferrimagnetism is a particular case of

antiferromagnetism in which the magnetic moments on the ‘A’ and ‘B’ sublattices are

in opposite direction having different magnitudes. Name ferrimagnetism is due to

Neel [7] who developed a general theory of the subject. High frequency ferrites was

initiated by the work done by J.L. Snoek [8] who found that high frequency range is

associated with these materials.

Fig. 1.6 Ferrimagnetism (arrows represent atomic magnetic dipoles) [5]

Fig. 1.7 Ferrimagnetic material with and without applied field

(arrows represent atomic magnetic dipoles)

Superparamagnetic materials


6

Superparamagnetic materials are those materials in which single-domain

ferromagnetic or ferrimagnetic particles show no long-range order between particles.

The shape of M-H curve is similar to paramagnetic materials except that the

magnetization in low to moderate field is much larger.

Fig. 1.8 Plot of coercivity as a function of particle diameter.

The particle becomes single domain at radius Dc

1.5 Types of ferrites with respect to their magnetic

properties

[5].

In case of nanocrystalline materials, domain size and crystal dimensions

become approximately comparable. In multidomain crystals, magnetization and

demagnetization is due to the rotation of domain walls but in case of single domain it

is harder to demagnetize the crystal by applying magnetic field. This is due to the

disruption of spin-spin coupling within domain. If the particle size is further reduced,

the number of spin decreases and the force aligning them becomes weaker. At last,

this force is too weak to overcome thermal randomization and in the absence of

applied magnetic field, the spins are randomly oriented. The crystal then becomes

superparamagnetic [5].

There are two types of ferrites with respect to their magnetic properties as:

1. Soft ferrites

2. Hard ferrites

Soft ferrites can be easily magnetized and demagnetized. These materials have

low coercivity and high saturation magnetization. Soft ferrites came into commercial

production in 1948. They consist of compound oxides consisting of iron oxide


7

(Fe2O3) together with other oxides such as Mn, Ni, Fe or Mg which have a chemical

complicated composition e.g NiO.Fe2O3 and F3O4 etc. In their final form, they are

usually brown colored ceramics. Their saturation magnetization is typically, Ms =

0.2x 106 A/m (Bs=0.25 T), with Hc=8 A/m (0.1 Oe) and maximum permeability µr

1.6 Types of ferrites with respect to their structures

=1500 [9]. The structure of these ferrites is cubic spinel which will be discussed in the

next section.

Hard ferrites are difficult to magnetize or demagnetize. These materials are

permanent magnets and have high coercivity and moderate saturation magnetization.

The most important hard ferrites are Alnico, barium ferrites and strontium ferrites [9].

Ferrites can be classified with respect to their structure as follows [10]:

1. Spinel cubic ferrites

2. Hexagonal ferrites

3. Garnets

Our total research work including thin films on spinel cubic ferrites, therefore

we discuss in detail only spinel cubic ferrites.

1.6.1 Spinel cubic ferrites This family of ferrites is most widely used in electronic industry. These ferrites

have very high electrical resistivity at room temperature except Fe3O4 and low value

of eddy current losses which make them versatile for their use at microwave

frequencies.

A large numbers of ferrites posses the structure of the natural spinel, MgAl2O4

which is a stable structure. The structure of the spinel was first determined by Bragg

and Nishikawa in 1915 [11]. Spinel ferrites are predominantly ionic. Many different

cation combinations may form a spinel structure, it is almost enough to combine any

three cations with a total charge of eight to balance the charge of anions. The limit of

the cation radii are approximately 0.4 to 0.9 Å (based on the oxide ion radii R0 of 1.4

Å). The important spinels from magnetic point of view are Fe2O3.


8

1.6.2 Spinel structure

Fig. 1.9 Location of A and B-sites in unit cell

Fig. 1.10 Local atomic arrangement for (a) tetrahedral site (b) octahedral site

in spinel structure [12]

The unit cell of the spinel structure (space group 3Fd m ) contains 8 formula

units in a cubic closed packed arrangement of the oxygen anions. The formula can be

written as A8B16O32

In octahedral sites, interstice is at the centre of an octahedron formed by 6

lattice anions. Four anions touching each other are in plane, the other two anions sites

. The anions are the largest and they form fcc lattice within these

lattices. Two types of interstitial position occur and these are occupied by metallic

cations. There are 96 interstitial sites in the unit cell, 64 tetrahedral (A) and 32

octahedral (B) sites.

Tetrahedral Site

In tetrahedral (A) site, the interstice is in the centre of a tetrahedron formed by four

lattice atoms. In this configuration, four anions are occupied at the four corners of a

cube and the cation occupies the body centre of the cube. For charge neutrality of the

system on 8 tetrahedral sites are occupied by cations out of 64 sites per unit cell in fcc

crystal structure.

Octahedral Site


9

in the symmetrical position above and below the centre of the plane formed by four

anions. For charge neutrality, 16 tetrahedral sites are occupied by cations out of 32

sites in a spinel structure.

Tetrahedral site has 12 nearest B-atoms and each B-atom in an octahedral site

has 6 nearest A-atoms, as shown in the fig. When A and B-atoms are both magnetic

elements, exchange interaction exist between A and B-atoms and the number of

nearest neighbor exchange interactions for each site should be also different. This

difference in number of exchange interactions depend on the crystallographic position

of each magnetic element and are important for magnetic properties of ferrites. Due to

this reason, magnetism in cubic spinel ferrites is strongly related to cation distribution

between tetrahedral and octahedral sites [12].

1.7 Types of spinel ferrites

There are three types of spinel ferrites due to their cations distribution on

tetrahedral (A) and octahedral (B) sites.

1. Normal spinel ferrites

2. Inverse spinel ferrites

3. Intermediate spinel ferrites

Normal Spinel Ferrites

The unit cell of spinel structure has 8MO.Fe2O3 molecules in 8 M2+ ions occupy 8

tetrahedral sites and 16 Fe3+

1 2 4( ) [ ]A site B siteD T D T Oδ δ δ δ− − − −

ions occupy 16 octahedral sites [13]. The general formula

for normal spinel is as follows:

Here D= divalent cations

T= trivalent cations

If δ =0, then the structure will be Normal.

Inverse Spinel Ferrites

In the inverse spinel structure, 8 M2+ ions occupy 8 octahedral sites and 16 Fe3+

1 2 4( ) [ ]A site B siteD T D T Oδ δ δ δ− − − −

ions

are divided in such a way that 8 occupy octahedral sites and 8 occupy tetrahedral sites

[14]. We can say that non-magnetic ions occupy 8 B-sites, whereas the iron is divided

between A and B sites. Here in the genral formula

If δ =1, the structure will be inverse.


10

Intermediate Spinel Ferrites

These are the ferrites having ionic distribution between normal and inverse spinel

ferrites are known as mixed ferrites.

e.g. 2 3 2 31 1 1 4( ) [ ]A BM Me M Me Oδ δ δ δ

+ + + +− − +

Here ‘δ ’ is called inversion factor and depends on the method of preparation and

nature of constituents of the ferrites. For mixed ferrites it is 1/3.

1.7.1 Substitutional ferrites Substituted ferrites are more complex than the normal ferrites. Sometimes

ferric ions are replaced by trivalent ions of another metal. Magnetization’s effect

depends on the site preferred by the substituent. It is difficult to accurately predict the

ion distribution in advance [13]. 1.8 Phases of Fe oxides Iron has total fifteen phases which have different physical, chemical and

magnetic properties [15]. The most important and distinct phases are FeO (wustite),

α- Fe2O3 γ (hematite), -Fe2O3 (maghemite) and Fe3O4 (magnetite). All phases are

composed of an O2-

Two forms of iron oxides are antiferromagnets: Fe

sublattice with iron ions occupying different interstitial sites. The

valence of the iron can be +2, +3 or mixture of two, depending on the type of iron

oxide. Magnetically, phases of iron oxides can be separated into two gourps:

antiferromagnets and ferrimagnets [16]

xO (wustite) and Fe2O3

(hematite). FexO the most reduced form of iron oxide, is an insulator that has a rock-

salt structure with a lattice parameter of 4.278 Å to 4.305 Å for 0.9<x<0.95 that

consists of Fe2+ as shown in the Fig.1.11 (a). It has a bulk Neel temperature of 198 K,

below which the magnetic moments align in ferromagnetic sheets along the [111]

direction. Each [111] sheet is antiferromagnetically aligned and the moments point

perpendicular to these sheets [17]. α-Fe2O3 is a fully oxidized phase of Fe oxide. It is

also insulating and has a corundum structure with a lattice constant of 5.424 Å,

consisting of only Fe3+ ions as shown in the fig.1.11 (b). It has a bulk Neel

temperature of 958 K with the moments aligned ferromagnetically in (111) planes.

Again each (111) sheet is antiferromagnetically aligned. The direction of the moments

relative to these sheets is temperature dependant. Below 260 K, these moments are


11

perpendicular to the (111) direction while above 260 K, they are parallel to the (111)

direction [18].

The other two phases of iron oxide, magnetite (Fe3O4 γ) and maghemite ( -

Fe2O3) are ferrimagnetic. Both have inverse spinel structure as shown in the Fig.1.11

(c) Maghemite is considered the fully oxidized version of magnetite, has only Fe3+

γ

ions that occupy both tetrahedral and octrahedral interstitial sites with an average of

1/6 of the octahedral sites per unit cell being vacant. The chemical formula of -

Fe2O3 can be written as [Fe3+]tet. [Fe3+5/3+V.S1/3] O4, showing unequal population of

both sites plus the additional vacancies (V.S) on octahedral sites. The moments on the

octahedral and tetrahedral sites are antiferromagnetically aligned, but due to unequal

population of each site, there is a residual net magnetism, pointing in the direction of

the octahedral moments. The magnetic easy axes for maghemite are also along the

[111] direction with a calculated Curie temperature of 948 K [19]. Bulk maghemite is

metastable, transforming into hematite at 670 K, but it has been shown to be

metastable in thin films at room temperature and was the basis of some early magnetic

storage media [20].

Fig. 1.11 Crystal structure for different phases of iron oxides (a) Wustite (b)

Hematite (c) Magnetite and Maghemite [16]


12

Magnetite is composed of both Fe2+ and Fe3+ ions. The Fe2+ ion sit at

octahedral sites while the Fe3+ ion occupy both octahedral and tetrahedral sites. At

room temperature the distribution of Fe3+ and Fe2+ in the octahedral site is presumably

random with only short-range order [21]. Again the magnetic moments

Fig. 1.12 Schematic representation of density of states (DOS) for

a) Normal non magnetic metal b) Normal ferromagnetic metal c) Half-

metallic ferromagnetic

on the octahedral sites are antiferromagnetically coupled to the moments on the

tetrahedral sites, with the magnetic easy axis along the [111] direction [22-23]. The

moments on the tetrahedral and octahedral Fe3+ ions cancel leaving the moments on

the Fe2+ ions uncompensated with results in the net magnetism. Fe3O4 has unusual

electronic properties. At room temperature, it is conducting while below Tv 120 K the

electrical conductivity drops by two orders of magnitude [24-25]. The Verwey

transition is a metal-insulator transition accompanied by a structural change from

cubic to monoclinic. Magnetite is considered as half-metallic ferrimagnet, where there

is 0.5 eV gap in the majority spin band at EF and exhibits normal metallic behavior

for the minority spin electrons as observed in several local density of state

calculations of the band structure of Fe3O4 [26-28].

From an itinerant electron point of view, the conductivity is a result of the

partially filled 3d band of the octahedral-site Fe atom and the Verwey transition is

argued to be a result of electron correlation and electron-phonon interactions that

cause a band splitting below Tv

From an ionic picture, the conductivity is due to the rapid hopping of minority

spin electrons between octahedral Fe

[29].

2+ and Fe3+ ions [30-31]. The Verwey transition

is then a result of an ordering of the Fe2+ and Fe3+ on the octahedral sites which

‘freezes out’ the hoping of electron between them. Fe3O4 is extensively used in a

large number of technological applications, particularly, in recording media industry


13

[32]. Furthermore it is of great interest in geology and archeology because of its

abundant occurrence in the earth’s crust and it impact on the local magnetic field.

1.9 Interactions in ferrimagnetics The interaction energy between the two atoms having spin ‘Si’ and ‘Sj

2 .e i jE J S S= −

’ was

shown by Heisenberg as follows [33]

Where ‘Je’ is the exchange intergral and this integral is a measure of the extent to

which the electronic charge distribution of two atoms concerned overlaps one another.

The two electrons under consideration spend a fraction of their time around the nuclei

of both atoms. The Pauli Exclusion Principle does not permit the two electrons with

the same spin to occupy the same energy state and the electrons must therefore be

‘exchanged’ between two atoms. This direct exchange interaction may be positive or

negative. The magnitude and sign of the exchange integral depends on the ratio of D/d,

where ‘D’ is the atomic or ionic separation of the interacting atoms or ions, and ‘d’ is

the diameter of the electron orbit concerned [34].

A study of the ionic arrangements within a spinel crystal shows that the direct

view, as it were, of metal ion is often obscured, partially or wholly, by an intervening

oxygen ion. Thus direct overlap of the electronic charge distributions of the cations is

not very probable. Mechanisms by which the negative interaction may be obtained,

and in which the oxygen ion plays an important role, have been suggested via

“superexchange” [35] and “double exchange” [36].

Superexchange

This type of indirect exchange normally extends from very short range-interaction to a

longer range. The idea of this exchange was given by Kramers in 1934 [37] and

theory was developed by Anderson in 1950 [38]. According to Kramers that this

exchange occurs through a non-magnetic atom. This exchange is important in ionic

solids such as transition metal oxides and fluorides, where the bonding orbitals are

formed by the 3d electrons in the magnetic transition metal atoms and the 2p valence

electrons in the diamagnetic oxygen or fluorine atoms. The size of the superexchange

depends on the magnitude of the magnetic moments on the metal atoms, the metal-

oxygen (M-O) orbital overlap and the M-O-M bond angle [39].


14

Double exchange

It is the indirect exchange in which the adjacent ions of parallel spins via

neighboring oxygen ions are considered. This arrangement requires the presence of

ions of the same metal in different valence states [40].

As an example in Fe3O4 which is overall ferrimagnetic but contains iron in two

different valence states Fe2+ (3d6) and Fe3+ (3d5) that are ferromagnetically coupled

[41]. The double exchange favors positive interactions. It will not account for

negative A-B interaction in ferrites.

The super-exchange and double exchange in Fe3O4 are shown in the following

fig. 1.13.

Fig.1.13 Crystallographic and magnetic structure in Fe3O4

1.10 Thin film ferrites

, near tetrahedral and

octahedrally (site A) and octahedrally (site B) coordinated Fe atoms. Here SE

represents “Superexchange” and DE represents “Double Exchange” [39].

Thin films of ferrites have technological importance as catalysts anticorrosives,

and magnetic devices. In particular, magnetite, as a half-metallic material, is an

attractive candidate for applications in spin electronics and magnetic recording.

In applications of magnetite in thin magnetic films the morphology of the

layers as well as the structure and composition of the surface are crucial factors

for the functionality [42].


15

1.11 Magnetic nanoparticles Generally nanosized object is a physical object that is different in properties

from the corresponding bulk materials and having maximum dimension of 100 nm.

Magnetic nanoparticles are mostly size dependent. Dimension of magnetic domain is

very small. If the magnetic particle dimension is comparable to magnetic domain

dimension, then the properties of those particles are very important in magnetism [43].

Nanotechnology is dealing with; single nano-objects and materials, devices

based on these materials and processes that take place in the nanometer range [43].

Nanomaterials are classified as follows;

• Nanostructured materials are materials isotropic in the macroscopic

composition and consisting of contacting nanometer-sized units as

repeating structural elements.

• Nanodispersed include a homogenous dispersion medium (vacuum, gas,

liquid and solid) and nanosized inclusion dispersed in this medium and

isolated from each other. Actually these are nanopowders whose grains are

separated by thin layers of light atoms which prevent them from

agglomeration [43].

1.12 Applications of ferrites Due to high electrical resistivity, these materials are very important in

electronic industry. The market price of these materials is very low compared to other

electroceramic: $33/kg for varistors, $330/kg for thermistors and $ 1100/kg for

ceramic capacitors [44].

For high-frequency applications the conductivity of metals limits their use and so we

must turn to magnetic insulators. These materials must of course exhibit the usual

properties associated with soft ferromagnets: high permeability, low coercivity and

high saturation magnetization. In these applications, soft ferrites are used widely.

Soft ferrites are also used in frequency selective circuits in electronic

equipment for example in telephone signal transmitter and receivers. Mn-Zn ferrites,

which are sold under the commercial name of ‘ferroxcube’, is widely used for

applications at high frequencies of up to 10 MHz, while beyond that frequency Ni-Zn

ferrites are preferred because they have lower conductivity. Another area where

ferrites find wide applications is in antenna for radio receivers. Almost all radio

receivers using amplitude modulation of signals are now provided with ferrite rod


16

antennae. Other applications include wave-guides and wave shaping for example in

pulse-compression systems.

The permeability of these materials does not change much with frequency up

to a critical frequency but then decays rapidly with increasing frequency. The critical

frequency of these materials varies between 10 MHz and 100 MHz. The saturation

magnetization of ferrites is typically 0.5 T, which is lower than iron and Co-alloys.

For high frequency applications, beyond 100 MHz, there are other materials

such as hexagonal ferrites which have special properties which make them suitable for

use at high frequencies. These materials are uniaxial with magnetic moments confined

to the hexagonal base plane [45].

1.13 Densities of ferrites The densities of ferrites are significantly lower than those of their thin metal

counterparts, thus, a component of the same size would be lighter in a ferrite.

However, this advantage can disappear due to low saturation. But density can be

increased by hot pressing. This resulting density is called sintered density [46].

1.14 Porosity in ferrites Porosity has effect on all the properties especially on mechanical and magnetic

properties. Porosity can be found as;

1 100bulk

x ray

P ρρ −

= −

Here ‘ P ’ is the porosity, ‘ bulkρ ’ is the bulk density and ‘ x rayρ − ’ is the x-ray density.

The range of porosity is 1% to 15% depending on the ferrites [46]. The ferrites having

porosity less than 1% are particularly valuable in the manufacture of devices such as

recording heads. High density ferrites having high permeability are used for

transformer cores. 1.15 Hardness of ferrites The importance feature for recording head is the ferrites high hardness; which

improves it wear resistance. The hardness of ferrite was measured on a limited

number of samples. The results in terms of the Vicker Pyramid number were 600 to

700 for Mn-Zn ferrites and 800 to 900 for Ni-Zn ferrites [46].


17

1.16 Dielectric behavior of ferrites Dielectric properties are most important in ferrites which depend on

preparation conditions e.g. sintering time and temperature, type and quality additives.

In order to improve the frequency performance (such as high quality factor, high

insulating resistivity, low dielectric constant and loss), many substitutions are possible.

We have also tried some substitutions which are described in chapter 4.

1.17 Electrical resistance of ferrites Due to high electrical resistivity, ferrites play a crucial role in electronic

industry. Electrical properties of ferrites depend on the following factors [34]:

• Chemical compositions

• Heat treatment during preparation

• Methods of preparation

On the basis of Verway hopping mechanism [47], we can explain the variation

of dc-electrical resistivity of ferrites. According to this mechanism, hopping of

electrons play a technical role in electrical conduction of ferrites between the ions of

same element but of different valence states present at octahedral sites [48]. Electrical

resistivity of ferrites decreases with the increase of temperature. This shows that

ferrites have semiconductor behavior as a function of temperature [49]. The typical

range of resistivity of ferrites from 10-2 Ω-cm to 1011 Ω-cm at room temperature

depending on chemical composition of the materials [46].

1.18 Magnetic behavior of ferrites The magnetization of ferrites can be discussed by comparing them with

ferromagnetic materials. Ferromagnetic materials have high magnetization (emu/cc),

even in polycrystalline form by the application of relatively small magnetic field.

In case of ferromagnetic materials, the individual atomic or ionic moments

arising from unpaired spins are permanent, and interact strongly with one another in a

manner which tends to cause parallel alignment of the nearby moments. The moments

of a large number of neighboring ions are thus parallel, even in the absence of an

applied field. These regions or domains, of spontaneous magnetization exist in both

single and polycrystalline materials, and within a domain the value of the saturation

magnetization M, is the maximum that can be achieved in the material at the given

temperature [50].


18

In case of ferrimagnetic materials, two sets of moments A and B exist. They

are aligned in opposite direction but with different magnitudes; so their effect is

partially cancelling and give us some overall net magnetization. The name

ferrimagnetism is due to Neel [51]. These materials are also known as ferrimagnetics,

consists of two sublattices formed by the tetrahedral and octahedral sites in the spinel

lattice.


19

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18. L. Neel, Rev. Mod. Phys. 25 (1953) 58.


20

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(1947) 181.

32. J.C. Mallinson, The foundation of magnetic recording, (1993).

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21

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Chapter 2 Literature Survey

22

Literature Survey Nanotechnology is the control of matter at dimensions of less than 100 nm and

size changes the properties. The physical and chemical properties of the nanomaterials

mostly depend on their size, shape and surface morphology. Physicists, chemists,

engineers and material scientists are focusing their concentration to develop simple

and effective methods to synthesize magnetic nanoparticles with controlled size and

shape. At nanoscale, structural, electrical and magnetic properties are not the same as

they are in bulk size. For example when the critical diameter of magnetic nanoparticle

lies in the range of few tens of nanometer, superparamagnetic behavior exists. Above

the superparamagnetic limit particle has single domain (without domain wall) and is

uniformly magnetized with all the spins aligned in the same direction as the applied

field. Since there is no domain wall at this stage, magnetization will be reversed by

spin rotation which increase the coercivity of small nanoparticles.

Ferrites are mainly composed of ferric oxide α-Fe2O3 and also called ceramic

like ferromagnetic materials. The saturation magnetization of ferrite is less than

ferromagnetic materials but have advantages, such as applicability at high frequency

and greater electrical resistance. Researchers have been engaged in the development

of new ferrites with improved manufacturing processes and properties. Manufacturing

of ferrites is little complicated as compared to ferromagnetic materials due to their

narrow range of stoichiometry and extra heat treatments. Today ferrites are used as

electronics parts, and therefore dimension must be exact and properties should be

uniform.

Due to different kinds of ferrites, research on ferrites is very vast; it is very

difficult to collect the whole information about all types of ferrites in every aspect. So

we are restricting ourselves to present a review of experimental facts that is related to

present study. The present literature review includes Cu-Zn ferrites, Al3+ doped

ferrites, Mn-Cu with Zn2+ and Ni2+ doped ferrties, Co-Mn nanoparticles and Fe3O4

2.1 Cu-Zn ferrites

thin films with pulsed laser deposition and useful comparison with sputtering and

MBE techniques.

Cu-Zn ferrites were discovered by Kato and Takei in 1932 and the name

designated for this material was oxide core. These materials have resistivity 103 Ω-m


23

and can be used for high frequency applications. These ferrites are sintered at least

1000 °C to reduce Cu2+ ion to Cu1+ ion [1] at high temperatures.

Rana et. al [2] have reported the effect of compositional variation on the Curie

temperature, magnetic moment and saturation magnetization of Cu1-xZnxFe2O4

system. They measured the Curie temperature from ac-magnetic susceptibility using

mutual induction technique. They reported that up to x=0.75, saturation magnetization

increased but further increase in Zn-content, decreasing trend is exhibited. Also Y-K

angle increases gradually with increasing Zn-content. They concluded that mixed Cu-

Zn ferrites exhibit a non-collinearity of the Y-K type.

Patil et. al. [3] have reported the magnetic properties of Cu1-xZnxFe2O4

ferrites (x=0, 1) by means of Mossbauer spectroscopy. They found that the systematic

dependence of the isomer shift, quadrupole interactions and nuclear magnetic fields

of 57Fe3+ ions in both A and B-sites as a function of Zn-content. Variation of nuclear

magnetic fields at A and B-sites are explained on the basis of A-B and B-B super

transferred hyperfine interactions.

Ravinder [4] studied the thermoelectric power studies of various compositions

from room temperature to well beyond the Curie temperature by different methods.

He confirmed that all compositions of ferrites show n-type semiconductors. The

charge carrier concentrations have been found through Seebeck coefficient.

In another work, Rana et. al [5] studied the effect of compositional variation

on porosity, magnetic properties and grain size of Cu1-xZnxFe2O4

Khalid Mujasam Batoo et. al [8] have studied the influence of Al doping on

electrical properties of Ni-Cd nano ferrties. The results obtained show that real and

imaginary parts of the dielectric constant, loss tangent and a.c conductivity shows

normal behavior with frequency. They explained the dielectric properties and a.c

conductivity of the samples Ni

(x=0.0, 0.25, 0.50,

0.75, 1.0) ferrites. From microstructural analysis, he follows that both porosity and

coercivity decreases with Zn-content. He concluded that coercivity is inversely

proportional to grain size with Curie temperature increased from 538 K to 560 K and

he related this decrease in coercivity as a function of grain size with Neel’s

mathematical model treating the demagnetizing influence of non-magnetic material in

mixed ferrites. Recently, Cu-Zn based ferrites have been synthesized, exhibiting high

Curie temperature with a little compromise on initial permeability [6-7].

0.2Cd0.3Fe2.5-xAlxO4 (0.0≤ x ≥0.5) on the basis of space


24

charge polarization according to Maxwell-Wagner two layer model and the Koop’s

phenomenological theory.

In another work Batoo et. al [9] have reported the electrical properties of Al

doped MnFe2O4 ferrites using ac impedance spectroscopy as a function of frequency

at different temperatures. In this work they reported that from impedance spectra it is

found that real and imaginary parts of the impedance decreases with increasing

frequency and both are found to decrease with Al doping up to 20% and then increase

with further increasing the Al concentration.

I.H. Gul et.al [10] have been synthesized CoFe2-xAlxO4 (for x=0.00, 0.25,

0.50) by sol-gel method. They studied the effect of Al3+ ions on structural, Curie

temperature, DC electrical resistivity and dielectric properties of CoFe2-xAlxO4

system. DC electrical resistivity has been explained by Verwey’s hopping mechanism.

They reported that activation energy increases with Al+3 ions and variation of

dielectric constant has been explained on the basis of space charge polarization.

Chhaya et. al [11] have investigated the NiAlxCrxFe2-2xO4 system for x=0.6 to

0.9 with a view to determining the effect of changing the Fe:Al:Cr ratio on the cation

distribution and magnetic ordering of the system with x-ray diffraction, magnetization

and Mossbauer effect measurements. In this work they reported that with increasing

concentration, the lattice parameter and saturation magnetization decrease but canting

angle increases indicating random canting of magnetic behavior.

Sam Jin Kim [12] studied Al substituted CoAlxFe1-xO4 (x=0.1, 0.2, 0.3, and

0.5) ferrites with x-ray, neutron diffraction, Mossbauer spectroscopy and vibrating

sample magnetometry. Temperature dependence of the magnetic hyperfine field

in 57Fe nuclei at A and B-site was analyzed based on the Neel theory of magnetism.

They reported that with increasing Al substitution the A-B and B-B interaction

decreased but A-A interaction increased and concluded that the reduction of magnetic

moment in 57Fe (A) and strength of A-A interaction are related to the covalency effect.

No previous work was found in the literature for the compositions Cu0.5Zn0.5Fe2-

xAlxO4

2.2 Co-Mn ferrites (for x= 0.0, 0.1, 0.2, 0.3, 0.4, 0.5) used in the present work.

Cobalt ferrites are promising materials for high density recording media

because of their high coercivity (Hc), moderate saturation magnetization (Ms),

chemical stability and mechanical hardness [13]. These materials are also good


25

candidates for the next generation magneto-optical recording media owing to their

large magneto-optical effect in visible wavelength range and good corrosion

resistance [14-16], compared to the widely used amorphous thin films [17-18]. There

are some hindrances in the applications and could be overcome. In polycrystalline

CoFe2O4, there is large media noise [19-20], which originates from light scattering at

grain boundaries and the irregular shape of the read-write domain (as a data bit). We

can lower the media-noise by reducing the grain size of polycrystalline oxides to the

nanoscale. Another limitation is their high Curie temperature of 520C which could not

meet the needs of the commercial disk writing temperature [21]. Metal substituted

cobalt ferrites are suitable for magneto-mechanical strain sensors and activators

applications [22]. Many researchers proposed Mn substituted cobalt ferrites in order

to tailor their magnetic and magneto-mechanical properties.

Kwang Joo Kim et. al [23] have investigated the effect of Mn doping on the

structural and magnetic properties of CoFe2O4 thin films prepared by sol-gel method.

They reported a large saturation magnetization for both MnxCo1-xFe2O4 and

MnyCoFe2-yO4 films compared to that of CoFe2O4. They explained such

enhancement of magnetization in terms of breaking of ferrimagnetic order induced by

the Co2+ migration.

B. Zhou et. al [24] prepared CoFe2-xMnxO4 (x=0-2.0) nanocrystalline thin

films and powders by sol-gel process. The reported that substitution of Mn+3 for Fe+3

causes the migration of Co2+ from A to B-site and finally lead to phase transformation

and decreases the saturation magnetization (Ms), Curie temperature (Tc) and

coercivity (Hc) with increasing Mn content x. They concluded that this decrease of

Ms and Tc is related to weakly magnetic properties of Mn3+, whereas the decrease of

Hc is related to decrease of Co content on the B-site.

K. Krieble et. al [25] have prepared a series of Co0.1MnxFe2-xO4

M.K. Shobana et. al [26] synthesized Co

samples and

studied them using Mossbauer spectroscopy. They reported that increase in Mn

content decreases the hyperfine field strength at both sites but at unequal rates and this

increase the distribution width. They concluded that this effect is due to the relative

strength of Fe-O-X superexchange (x=Fe, Co, or Mn) and the different numbers of the

next nearest neighbors of A and B-sites.

0.5Mn0.5Fe2O4 nanoparticles using

sol-gel combustion method in which citric acid was used as the complexing agent.

They reported the structural, thermal and magnetic properties as a function of


26

calcination temperature. They concluded that saturation magnetization increases with

the increase of size of particles and the coercivity of the nanoparticles varies

significantly as calcination temperature increases. Also they reported that as-burnt

powder shows amorphous behavior and the spinel structure starts to appear at 500°C.

In this present work I have selected the same composition as reported by

Shobana et. al [26] and prepared by sol-gel combustion method. I obtained crystalline

behavior of as burnt powder. No previous report was found for such a low

temperature synthesis of Co0.5Mn0.5Fe2O4

2.3 Mn-Cu ferrites ferrites.

Cu substituted MnFe2O4 ferrites have been prepared by ceramic method after

considering the composition Mn1-xCuxFe2O4 (x=0,0, 0.25, 0.50, 0.75, 1.0) [27]. In

this report, the authors studied the microstructure analysis and discuss the porosity

with ‘Cu’ concentration. They reported that that porosity increases with Cu

concentration whereas coercivity increases up to x=0.50 and decrease of coercivity

after x=0.50 and was explained on the basis of Neel’s mathematical model treating the

demagnetizing influence of non-magnetic material in cubic crystals. Also they related

the decrease of coercivity with grain size to inter-granular domain wall movement

because of large porosity.

Rana et. al [28] have studied the effect of compositional variation on magnetic

susceptibility, saturation magnetization (Ms), Curie temperature (Tc) and magnetic

moments (µB). In this work they reported that by increasing ‘Cu’ concentration up to

x=0.50, Ms increases while Curie temperature decreases. After x=0.50, Ms decreases

while Curie temperature continue to decrease. This effect was explained by partially

related to the low magnetic moment of Cu2+ ions. Also Y-K angle increases gradually

with increasing Cu contents and extrapolates to 90° for CuFe2O4. They concluded

that mixed copper ferrites exhibit a non-collinearity of the Y-K type while MnFe2O4

Z. Zrmsa et. al [30] discussed magnetic bubbles in spinel ferrites films

considering the composition Mn

shows a Neel type of ordering.

In another work, Rana et. al [29] have studied the cation distribution in Cu-

substituted manganese ferrites, aiming to study the relationship between structural

parameters and concentration of the substituted copper ions. They concluded that

these ferrites belong to the family of mixed or partially inverse spinel.

0.5Cu0.5Fe2O4 and Ni-ferrties. These films were


27

prepared with chemical transport deposition technique and studied the magnetic

domains. Bitter technique was used to study magnetic domains by applying 3000 Oe

field. From the geometry of domain structure, saturation magnetization, characteristic

material length and domain wall energy were determined. Stability of cylindrical

domains was confirmed by ferromagnetic resonance. It was concluded that that the

behavior of bubbles in magnetic field with these films differ from that in garnets and

orthoferrites.

In our present work, we have synthesized Mn0.5Cu0.5-xZnxFe2O4 and

Mn0.5Cu0.5-xNixFe2O4

2.4 Fe

ferrites with sol-gel combustion technique and studied their

structural and magnetic properties. Alex Goldman [31] has reported these types of

materials as square loop ferrites but depending on the preparation conditions.

3O4

Half-metallic materials (majority-spin electrons were metallic, whereas

minority- spin electrons should be semiconducting with 100% spin polarization (i.e.

every mobile electron in the contact material has the same electron spin orientation)

of the charge carriers at the Fermi level are very attractive as potential applications in

spin electronics [32]. Fe

thin films

3O4 is a predicted half-metallic material [33-34]. Also it is

promising candidate for potential applications in spin electronics due to its high Curie

temperature (860 K) as compared to that of other half-metallic materials. Up to now,

many researchers have been prepared epitaxial or polycrystalline Fe3O4 thin films

with different deposition techniques such as molecular beam epitaxy (MBE) [35-36]

electron beam ablation [37], pulsed laser deposition from α-Fe2O3 γ, -Fe2O3 as a

target [38-41] and sputtering from an iron target [42-44]. Fe3O4 has a narrow range of

stoichiometry therefore many other phases like FeO, α-Fe2O3 γ and -Fe2O3 coexist

according to the specific deposition conditions [45-47]. However it is still difficult to

grow them with well defined compositions with pulsed laser deposition (PLD) from

different targets.

Tiwari et. al [48] have prepared (111) oriented Fe3O4 thin films independent

of substrate orientation with pulsed laser deposition technique on Si substrates of

different orientations, (111), (100) and (110). Single phase Fe3O4 was confirmed with

Raman spectroscopy. They suggested that all the films show ferromagnetic behavior

with saturation magnetization close to that of single crystal.


28

Parakash et. al [50] deposited Fe3O4 on GaAs(100) substrate by pulsed laser

deposition technique. In another work, Tiwari et. al [49] deposited Fe3O4 thin films

from α-Fe2O3 as a target with pulsed laser deposition on different substrates Si(111),

GaAs(100), Al2O3(001) and amorphous float glass without any buffer layer at a

substrate temperature of 450°C. XRD results show highly (111) oriented growth and

single phase nature of Fe3O4. They concluded that this highly (111) oriented growth

is due to huge lattice mismatch of substrates with Fe3O4.also reported (111) oriented

growth of Fe3O4 thin films and proved Verwey transition temperature at 122 K. All

their films show room temperature ferromagnetic behavior and saturation

magnetization close to the single crystal.

Kennedy et. al [51] have grown Fe3O4 films by laser ablation on Si(100) and

GaAs substrates after adjusting the substrate temperature of 450°C in an oxygen

atmosphere of 10-4 torr. In this work they used an epitaxial buffer layer of MgO (10

Å) thin and obtained (100) oriented Fe3O4 thin films. X-ray pole figure measurements

on these films indicate both Fe3O4 and MgO films are oriented cube on cube on the Si

and GaAs substrates. They concluded that hysteresis loops for the (111) and (100)

oriented films are very similar indicating the epitaxy has not significantly improved

the magnetic properties of the Fe3O4 thin films. Furthermore, they obtained high-field

magnetization values of 600-800 emu/cc for all the Fe3O4

Parames at. al [52] ablated Fe

films deposited on Si and

GaAs with and without MgO buffer layer. The reason for this high saturation

magnetization is due to the formation of Fe-rich regions within the films which would

significantly increase the magnetization. They concluded that these regions show

amorphous behavior and no iron peak exist in the x-ray diffraction pattern. They also

suggested that stoichiometry could indeed be the reason for the increase in

magnetization.

3O4 target with pulsed laser deposition technique

and deposited on Si(100) substrates in reactive atmosphere of O2 and/or Ar, with

different oxygen partial pressures. Their results show that a background mixture of

oxygen and argon improves the Fe:O ratio in the films as long as the oxygen partial

pressure is maintained in the range of 10-2 Pa range. They obtained a single phase

polycrystalline magnetite Fe2.99O4 at 483 K and working pressure of 7.8 x 10-2 Pa

with a saturation magnetization 490 emu/cc and Verwey transition temperature of 112

K close to the values reported in the literature for bulk material. They concluded that


29

stoichiometric magnetite with very good magnetic properties can be obtained at the

oxygen partial pressure of 4.7 x 10-2 Pa.

In another work, Parames et. al [53] have deposited magnetite on Si(100),

GaAs(100) and Al2O3(0001) at substrate temperature varying from 473 to 673 K by

pulsed laser deposition method in a reactive atmosphere of oxygen and argon, at

working pressure of 8 x 10-2 Pa. The influence of substrate on stoichiometry,

microstructure and magnetic properties were determined by XRD, conversion electron

Mossbauer spectroscopy (CEMS) and magnetic measurements. They concluded that

magnetite crystallite with stoichiometry varying from Fe2.95O4 to Fe2.99O4 are

randomly oriented on Si(100) and GaAs(100) and exhibit (111) oriented texture if

grown on to Al2O3(0001). Interfacial Fe3+ diffusion was present in both Al2O3 (0001)

and GaAs(100) substrates.

There are few reports about the annealing effect on the structural and magnetic

properties of Fe3O4 thin films on Si(100) exist [54-55], so there is the need to study

the effect of annealing temperature and time on the structural and magnetic properties

of Fe3O4 thin films on Si(100) substrates. This study is one focus of the present work.


30

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Chapter 3 Experimental Techniques

33

Experimental Techniques

3.1-Preparation Methods

A series of ferrites of various compositions and thin films were prepared with the

following methods;

1. Solid state reaction method

2. Sol-gel auto-combustion method

3. Pulsed laser deposition technique

The various prepared ferrites with their concentrations and techniques are given in

the Table (3.1).

Compositions Concentration/heati

ng temperature

Preparation

techniques

Characterization

techniques

Cu0.5Zn0.5Fe2-xAlxO

x= 0.0, 0.1, 0.2, 0.3,

0.4, 0.5 at 1100 °C

for 44 hrs. 4

Solid state

reaction

XRD, VSM, SEM

and Electrical

properties

Mn0.5Cu0.5-xNixFe2Ox= 0.0, 0.1, 0.2, 0.3,

0.4, 0.5 at 300 °C 4

Sol-gel

combustion XRD and VSM

Mn0.5Cu0.5-xZnxFe2Ox= 0.0, 0.1, 0.2, 0.3,

0.4, 0.5 at 300 °C 4

Sol-gel


Mn0.5Co0.5Fe2O500, 600, 700, 800

and 900 °C for 1 hr. 4

Sol-gel


Fe3O

As deposited on

Si(100) substrates

and annealed at 300,

400 and 450 °C

4 Pulsed laser

deposition XRD and VSM

Table: 3.1 Compositions with their concentrations/calcination temperature and preparation techniques


34

3.1.1 Solid State Reaction Method

Ceramic or solid state reaction method is the most common and one of the

simplest ways of preparing solids. It consists of heating together two non-volatile

solids which react to form the required product. This method is commonly used in

both industry as well as in the laboratory, and considered a best way to synthesized

oxide materials. The first high temperature superconductors were made by this

method.

The simple procedure is to take stochiometric amounts of oxides, grind them

in a pestle and mortar to give a uniform small particle size and then heat in a furnace

for several hours in a ceramic, alumina or platinum crucible according to the required

temperatures. Although this method is being used widely but it has several

disadvantages. Due to high temperature, a large amount of input energy is needed.

This is because the coordination numbers in binary or ternary compounds are high,

and it takes a lot of energy to overcome the lattice energy so that a cation can leave its

position in the lattice and diffuse to different sites. Sometimes the phase or

compounds may be unstable or decompose at such high temperatures [1].

Fig. 3.1 Pestles and mortars for fine grinding [1]


35

Fig. 3.2 Selection of porcelain, alumina or platinum crucibles [1]

3.1.2 Sol-gel combustion method Chemists, physists, engineers and material scientists have to face challenges to

synthesize oxide materials with exact structures, composition and properties. It is

difficult to control diffusion of atoms and ionic species through reactants and products

by ceramic method. Many attempts have been made to eliminate the diffusion control

problems of solids with different techniques [2]. One such successful technique is sol-

gel combustion technique. The term ‘combustion’ covers flaming (gas-phase),

smoldering (heterogeneous) as well as explosive reactions. The Combustion method

has been successfully used in the preparation of a number of magnetic, dielectric,

insulators and semiconductor materials. Some other advantages of “combustion

synthesis” are as follows [3]:

1. need simple equipment

2. formation of high purity product

3. produce nano particles

4. stabilization of metastable phases

5. formation of virtually any size and shape products

Fig. 3.3 Combustion process [3]


36

3.1.3 Pulsed laser deposition technique Pulsed laser deposition (PLD) was invented after the demonstration of ruby

laser in 1960. Smith and Tunner [4] in 1965 put forward new idea that thin films can

be deposited with the help of intense laser radiation. In PLD system, an intense laser

beam having energy density of a few J/cm2

1- Laser source is located outside the vacuum chamber and focused through a

quartz window

is focused onto the target, a portion of it is

absorbed and another reflected. Above a critical value of power density, material is

ejected from the thin surface region creating a vapor plume extending along the

direction normal to the target surface. The required power density depends on the

target material and laser pulse duration and wavelength. The material then travels

towards the substrate and deposited there. A schematic diagram thin film deposition

under PLD-system is shown in the fig. 3.4.

This technique is very much popular for the deposition of superconductors [5-

8] and diamond like carbons [9-16]. PLD-system has following key advantages [17],

2- Lasers are clean thermal sources that introduce minimal contaminations

3- Any kind of materials can be ablated due to high power densities.

4- Pulsed nature of the process gives precise control of the amount of

deposited material and produced better stoichiometry and properties of the

target than ordinary evaporation.

5- A standard PLD-system is cheaper than an MBE-system.

Fig. 3.4 Thin film deposition under PLD-system


37

The main drawback of PLD-system for commercial application is the

production of unwanted micron sized particles ejected during ablation process. These

can cause defects and change the properties of thin films. Also we can not deposit

large surface area with PLD-system.

The system which was used in the present study to deposit thin films on Si(100)

substrates has the following specifications [18]:

Table 3.2 Pulsed Nd:YAG Laser NL303 (EKSPLA) Specifications

Parameters Standard Specifications

Pulse duration, ns 3-6

Jitter (optical pulse to sync

pulse, standard deviation),

ns

0.5

Pulse energy, mJ

at 1064 nm

at 532 nm

at 355 nm

At 266 nm

800

360

240

80

Pulse energy stability

(standard deviation)

at 1064 nm

at 532 nm

at 355 nm

At 266 nm

1%

1.5%

3%

3.5%

Repetition rate, Hz 10 (and 10/N)

Near field intensity profile Hat top

Beam divergence, mrad <0.5

Pointing stability, µrad <50

Polarization Vertical / horizontal

Power consumption, KVA

(220 V AC; 50 Hz)

≤ 2.5

Dimensions, mm

Laser head

Power supply cabinet

455 x 120 x 120

330 (W) x 520 (d) x 670 (h)


38

3.2 Characterization Techniques To investigate the properties of ferrites, some precise techniques have been used

in our studies. The materials discussed in this thesis are characterized by the following

techniques;

• X-rays diffraction (XRD)

• Measurement of bulk density

• Scanning Electron Microscopy (SEM)

• Vibrating Sample Magnetometer (VSM)

• Dielectric properties measurement

• Electrical properties measurement

The brief description of these techniques is given below;

3.2.1 X-ray diffraction (XRD) X-ray diffraction is popular, versatile and non-destructive technique to study

the crystalline materials. Since crystal lattice has three-dimensional distribution of

atoms, arranged in such a way that they form a series of parallel planes separated by a

distance‘d’ which is called “inter-atomic distance”. For any crystal, planes are found

in different orientations each with its own specific d-spacing.

The wavelength of X-rays lies in the range of 0.5 to 2.5 Å which is

comparable to inter-atomic spacing in solids [19]. There are two ways to produce x-

rays, either by the deceleration of fast moving electrons in the metal target

(continuous spectrum) or by the inelastic excitation of the core electrons in the atoms

of target (characteristic x-rays) [ 20].

Diffraction in crystals occurs only when Bragg’s law is satisfied. This law

states that when radiation falls on a series of parallel planes equally spaced at a

distance ‘d’, then the path difference is 2dSinθ for the reflected rays, where ‘θ’ is

measured from the plane. The mathematical form of Bragg’s law [21] is

2dSinθ=mλ where m= 1, 2, 3,………


39

Fig. 3.5 Geometrical description of Bragg’s law

When interference occurs from many rows, then the constructive interference

peaks become very sharp with mostly destructive interference in between them. This

sharpening of the peaks as the number of rows increases is similar to the sharpening

of the diffraction peaks from a diffraction grating as the number of slits increases.

We can satisfy Bragg’s law either by varying ‘λ’ or ‘θ’ during experiment.

The methods in which these quantities can be varied as follows [22]:

Diffraction Methods Wavelength (λ) (Å)

Angle (θ) (degree)

Powder diffraction fixed variable

Rotating crystal fixed variable

Laue diffration variable fixed

Table 3.3 Diffraction method with their wavelengths and angle

We discuss only powder method because our samples are in the form of

powder except thin films. 3.2.2 Powder Method This is the most reliable method to find crystal structure and estimate the

crystallite size in the powdered specimen [23]. In this method a very fine powder is


40

put in a plate of glass or aluminum. Then a beam of monochromatic x-rays is incident

on the specimen. Each particle of the powder is in a tiny crystal oriented at random

with respect to the incident beam. Fig. 3.5 shows only one plane and diffracted beam

is formed. If this plane is rotated about the incident beam in such a way that then

reflected beam will travel over the surface of cone as shown in fig. This rotation does

not actually occur in the powder method, but the presence of large number of crystal

particles having all possible orientations is equivalent to this orientation, as some of

particles satisfy Bragg’s law with the incident beam. From the measured position of a

given diffraction line on the film, knowing ‘θ’ and ‘λ’, we can calculate the

interplanar distance‘d’.

In diffractometer method, for a given wavelength ‘λ’, the diffraction from a

plane (hkl) occurs at a certain angle ‘θ’. X-ray detector observed this x-ray beam. The

first x-ray diffractometer was used by Bragg for crystallography [24].

Powder x-ray diffractometer (Rigaku-D-MaxII-A) facility was availed in the

“Centre for Solid State Physics” University of the Punjab, Lahore, Pakistan but thin

film x-ray diffractometer (Rigaku) facility was availed in “Korean Advanced Institute

of Science and Technology” South Korea.

3.2.3 Measurement of bulk density The bulk density of complex shape is determined by “Hydrostatic Method”

which is a simple method based on “Archimedes Principle”. If the sample under study

has weight ‘Wa’ in air and ‘Wf

a

a f

WW W

ρ =−

’ in fluid i.e, water the density of the sample is given

by [25];

Here the difference in weight of the part in air compared to its weight suspended in

water permits the calculation of the density.

In the present work a densitometer (Gibitre-Instruement, Model: Electronic

Balance with serial number EBC2005081) was used for density measurement at

PCSIR Laboratories Paksistan. This instrument has the ability to determine the

density of the sample by comparing the weight obtained in air and with reference

liquid of known density using a precision balance.


41

3.2.4 Scanning Electron Microscopy (SEM)

Fig. 3.6 Schematic diagram of scanning electron microscopy [26]

It is a type of electron microscope that is used to produce high resolution three

dimensional images of a specimen surface. This technique is useful for looking at

particle or grain size, crystal morphology, magnetic domains, porosity surface defects.

Electrons are thermionically emitted from the cathode surface made of

tungsten or LaB6

An ordinary microscope can magnify the image of an object up to 1200 times

whereas the electron microscope can magnify the image up to 200,000 times. This large

magnification is due to the fact that the wavelength of a high-speed electron is much

and are accelerated towards the anode. Tungsten is used as it has the

highest melting point and lowest vapor pressure of all the metals, thereby favorable

for electron emission at highest temperature. The energy of the electron beam ranges

from a few spot size of 1 nm to 5 nm.

When the primary electron beam strikes the sample surface, the electrons

loose their energy by repeated scattering and absorption with the specimen and the

beam extends from less than 100 nm to around 5 µm on the specimen surface. The

energy exchanged between the electron beam and sample under observation, results in

the emission of electrons and electromagnetic radiations which are used to produce an

image. The resolution of scanning electron microscope ranges from 1 nm to 20 nm

[27].


42

lower than that of visible light, and so much higher resolution is possible. One major

advantage of SEM in comparison to TEM is the ease of specimen preparation, a result

of the fact that the specimen does not have to be made thin. In fact many conducting

specimens require no special preparation before examination in the SEM. On the

other hand, specimens of insulating materials do not provide a path to ground for the

specimen current and may undergo electrostatic charging when exposed to the

electron probe. This current can be of either sign, depending on the values of the

backscattering coefficient and secondary electron yield. Therefore, the local charge on

the specimen can be positive or negative. Negative charge presents a more serious

problem, as it repels the incident electrons and deflects the scanning probe, resulting

in image distortion or fluctuations in image intensity.

One solution to the charging problem is to coat the surface of the SEM

specimen with a thin film of metal or conducting carbon. This is done in vacuum,

using the evaporation or sublimation technique. When coating is undesirable or

difficult (specimen is very rough), specimen charging can often be avoided by

carefully choosing the SEM accelerating voltage [28]. 3.2.5 Vibrating Sample Magnetometer (VSM)

Vibrating sample magnetometer (VSM) determines the difference in magnetic

induction between a region of space with and without the specimen. Therefore, it

gives us the direct measurement of magnetization. If the sample is very short, then it

is difficult to measure the magnetization curve due to the demagnetization effects

associated with the short specimen. This method is good for the determination of

saturation magnetization (Ms).

To obtain M-H loops for ceramic and sol-gel prepared samples, we used

“Lake Shore7407” VSM at “Centre for Solid State Physics”, Punjab University

Lahore, Pakistan. In case of Fe3O4 thin films, we used “Ricken Denshi” VSM at

“Korean Advanced Institute of Science and Technology” (KAIST) South Korea.


43

Fig. 3.7 Schematic diagram of Vibrating Sample Magnetometer [29]

3.2.6 Dielectric Properties Measurement Most of the ferrites except Fe3O4 are very good dielectric materials and have

many technological applications ranging from microwave to radio frequencies. A

particular application requiring soft ferrites that has rapidly grown in importance in

the last few years is power supplies for computers, peripherals and small instruments.

A compact and efficient power unit can be obtained by using a technique known as

switched mode power supply (SMPS). In this technique, involving a dc to dc

conversion, one of the key elements is a high frequency transformer [29]. Hence it is

important to study the dielectric behavior of ferrites at different frequencies. The

dielectric constant є/

/

0

cdA

εε

=

measurements were carried out in the frequency range from 100

Hz to 1 MHz at room temperature using LCR meter (WK LCR 4275). The dielectric

constant was calculated by the following formula;


44

Here ‘C’ is the capacitance,‘d’ is the thickness, ‘A’ is the area of the sample and ‘ 0ε ’

is the permittivity of the free space.

The imaginary part of the dielectric constant ( //ε ) is a measure of the

absorption of energy by the dielectric from the alternating field. The dielectric loss

factor can be calculated by using the following relation;

1tan2 p pR C

δπ

=

Here ‘ δ ’ is the loss angle, ‘f’ is the frequency, ‘Rp’ is the equivalent parallel

resistance and ‘Cp

(tan )δ

’ is the equivalent parallel capacitance. The dielectric loss is also

measured in terms of loss tangent defined by the relation; // / tanε ε δ=

3.2.7 Electrical properties measurement Ferrites have high DC-electrical resistivity at room temperature. These

properties depend on chemical, compositions, various heat treatments and methods of

preparation [30-33]. DC-electrical resistivity of ferrites decreases with increasing

temperature showing the semiconductor behavior [34]. Resistivity of ferrites at room

temperature depends on chemical compositions [35]. The main cause of low

resistivity in ferrites is due to the simultaneous presence of ferrous and ferric ions on

equivalent lattice site (Octahedral). Resistivity of ferrites can be controlled by cation

distribution in B-site [36].

Two methods are used to measure high resistance:

a) constant voltage method

b) constant current method

Constant current method is used commonly for superconductors with four probe

technique. The resistivity of the semiconducting material is often measured with two

probe technique. This technique involves bringing two probes in contact with a

material of unknown resistance. The conduction electrons have resistance due to the

following reasons [37];

• The vibration of lattice ions due to the increase in temperature is major

cause of resistance. This temperature dependent resistivity is called

thermal resistivity.

• Imperfection and dislocations in a crystal is also a source of resistivity

which is lower than that due to lattice vibration.


45

Ferrites show semiconductor behavior with the rise of temperature. Resistivity

decreases according to Arrhenius equation [38];

0

EkTeρ ρ=

Here ‘ k ’ is the Boltzmann constant, ‘T’ is the temperature measured in Kelvin and

‘E’ is the activation energy which is the energy needed to release an electron from the

ion for a jump to neighboring ion, so giving rise to the electrical conductivity.


46

References 1. Lesely E. Smart, Elaine A. Moore, Solid State Chemistry, 3rd

2. Rao. CNR, “Combustion Synthesis”, In chemical approaches to the synthesis

of inorganic materials, New Delhi, Wiley Eastern Limited, (1994) 28.

. Edition, CRC

Press, (2005).

3. Kashinath C. Patil, S.T. Auna, Tanu Mimani, Current opinion in Solid State

and Materials Science 6 (2002) 507.

4. D.B. Chrisey, Graham K. Hubler, Pulsed laser deposition of thin films, Wiley

and Sons, INC, (1994).

5. Li, Q.et al., IEEE Trans. Appl. Superconduct. 5 (1995) 1513.

6. H.U. Habermeier, Appl. Surface Sci., 69 (1933) 204.

7. H.U. Habermeier, H. U., Mater. Sci. Eng. B-Solid State Mater. Adv. Technol.

13, (1992) 1.

8. H.U. Habermeier, Phys. C, 180 (1991) 17.

9. H.Minami, D. Manage, Y.Y. Tsui, R. Fedosejevs, M. Malac and R. Egerton,

Appl. Phys. A-Mater., (2001) 531.

10. A.A. Voevodin, M.S. Donley, and J.S. Zabinski, Sur. Coat Technol. 92 (1997)

42.

11. Q.R. Hou, and J. Gao, J. Phys. Cond. Matter. 9 (1997) 10333.

12. A.L. Karuzskii, N. N. Melnik, V. N. Murzin, V. S. Nozdrin, A. V.

Perestoronin, N. A. Volchkov and B. G. Zhurkin, Appl. Surface Sci., 92

(1996) 457.

13. A.A. Voevodin, , M.S. Donley, Surf. Coat. Technol. 82 (1996) 199.

14. A.A. Voevodin, and S.J.P Laube, Surf. Coat. Technol. 77 (1995) 670.

15. A.A. Voevodin, M.S. Donley, J.S. Zabinski, and J.E. Bultman, Surf. Coat.

Technol., 77 (1995) 534.

16. D.L. Pappas, L.L. Saenger, J.Bruley, W. Krakow, T. Gu and R.W. Collins. J.

Appl. Phys. 71 (1992) 5675.

17. J.R. Groza, James F. Shackelford, Enrique J. Lavernia, Michael T. Powers,

Materials processing Handbook, CRC Press, New York. (2007).

18. www.ekspla.com, “Plused Nd:YAG Laser NL303, Technical Description and

user manual, Vilnius, (2005).

http://www.ekspla.com/


47

19. Lesely E. Smart, Elaine A. Moore, “Solid State Chemistry”, 3rd

20. Rao. CNR, “Combustion Synthesis”, In chemical approaches to the synthesis

of inorganic materials, New Delhi, Wiley Eastern Ltd., (1994) 28.

.

Edition,Taylor and Francis, CRC press. (2005).

21. Kashinath C. Patil, S.T. Auna, Tanu Mimani, “Current opinion in Solid State

and Materials Science 6 (2002) 507.

22. D.B. Chrisey, Graham K. Hubler, Pulsed laser deposition of thin films, Wiley

and Sons, INC, (1994).

23. K.L. Horovitz, V.A. Johnson, “Solid State Physics, Vol.6, Academic Press,

New York and London, (1959).

24. B.D. Cullity, Introduction to magnetic materials, Addison Wesley Publishing

Companym (1972).

25. Ray F. Egerton, “Physical principles of electron microscopy, Springer Science

and Business Media Inc., (2005).

26. M. Ajmal, PhD-Thesis, Q.A.U. Isalamabad, (2008) 41.

27. Ray F. Egerton, “Physical principles of electron microscopy, Springer Science

and Business Media Inc., (2005).

28. B.D. Cullity, Introduction to magnetic materials, Addison Wesley Publishing

Companym (1972).

29. Raul Valenzuela, Magnetic Ceramics, Cambridge University Press, (1994).

30. T. Abbas, M.U. Islam, M.A. Chaudary, Mod. Phys. Lett. B, 9 (1995) 1419.

31. A.J. Deckar, Solid State Physics, The Mc Millan Press Ltd. London, (1995).

32. J. Smit, H.P.J. Wijn, Ferrites, Jhon Wiley and Sons, New York (1959).

33. E.J. Verwey, Heilman, J. Chem. Phys. 15 (1947) 174.

34. M.U. Rana, M.U. Islam, T. Abbas, K. Turkish, J. Phys., 19 (1995) 1137.

35. J.B. Goodenough, A.L. Loeb, Phy. 98 (1955) 391.

36. M.A. Semary, M.A. Ahmed, Y. Abbas, J. Mater. Sci., 18 (1983) 2890.

37. G.T. Meaden, Electrical Resistance of Metals, Haywood Books, London,

(1966).

38. E.J. Verwey, Heilman, J. Chem. Phys. 15 (1947) 174.

Chapter 4 Structural, magnetic and electrical properties of Al3+

substituted CuZn-ferrites

48

Structural, magnetic and electrical properties of Al3+

4.1 Motivation


(This work has published in Chinese Journal of chemical Physics, 2010)

Ferrite materials have attracted a considerable attention of the researchers

through decades due to their interesting soft magnetic properties and high frequency

applications [1]. A proper choice of cations along with Fe2+, Fe3+ ions and their

distribution between tetraherdral (A-site) and octahedral (B-site) sites of the spinel

lattice, imparts useful and interesting electrical and magnetic properties to the spinel

ferrites. Further tailoring of these properties using appropriate methods of preparation,

chemical composition, sintering time and doping additives always help to improve the

technological applicability of the ferrite materials [2]. It is essential to control the

electrical resistivity of the spinel ferrites in order to corporate these materials for a

wide range of applications. This can be achieved in two ways: controlling the

sintering temperature and by proper elemental substitution. Excellent dielectric

properties of ferrites further extend their application range from microwave to radio

frequencies. The useful frequency range is fixed by the onset of resonance

phenomenon for which either the permeability starts to decrease at a critical frequency

or the losses rise rapidly [3]. Recently, Cu-Zn based ferrites have been synthesized,

exhibiting high Curie temperature with a little compromise on initial permeability

[3,4]. The presence of Cu ions in ferrites activates the sintering process leading to

increase in density and decrease in losses. While, it is well known that Zn content

exerts important influence on the microstructure and hence on the magnetic properties

of ferrites. The substitution of Al3+ in ferrites could lower the dielectric constants that

warrant their applications for high frequency applications, for instance as micro wave

absorbers.

In the present work, we have systematically investigated the effect of Al3+ ion

substitution on the structural, magnetic and electrical properties of Cu0.5Zn0.5Fe2O4

.

The electrical behavior of the samples have been discussed in context of temperature

dependent resistivity, and frequency dependent dielectric constant (ε′), tangent of

dielectric loss angle (tanδ), and dielectric loss factor (ε).



49

4.2 Sample Preparation Samples of Cu0.5Zn0.5AlxFe2-xO4 (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5) ferrites have been

prepared by the standard solid state reaction technique using analytical grade reagents.

Low cast CuO (99%), ZnO (99%), and Fe2O3 (97%) in their respective stoichiometric

ratios were mixed to prepare the ferrite samples. Grinding of every sample with

specific composition was carried out in agate mortar and pestle for 4 hours. The

samples were calcined in the muffle furnace at 800 C for 8 hours. After the in-situ

cooling of the samples in the furnace, each sample was ground again for 2 hour. The

samples in powder form were pelletized (dia-15 mm) using Apex hydraulic press by

exerting a uniaxial pressure of 5 tons for 3 minutes. The samples were annealed at

1100 ˚C for 44 hours in order to get the required phase.

The investigation of the crystal structure was carried out using a Rigaku D-Max

II-A, diffractometer system with Cu Kα (λ = 1.5406 Å) radiation. Surface morphology

and microstructural features such as grain size and porosity were examined using

Hitachi S-3400, scanning electron microscopy (SEM). The grain size was measured

by using the line intercept method.

As ferrites are highly resistive materials, therefore two probe method was

employed to determine the electrical resistivity of the samples in the temperature

range from room temperature (RT) to 480 K. Frequency dependent (up to 1 MHz) RT

measurements of dielectric constant and dielectric loss were obtained using a

QuadTech-1920 LCR Meter. Magnetic characterizations were performed using a Lake

Shore-7404 vibrating sample magnetometer (VSM).



50

4.3 Results and discussion

Fig. 4.1 XRD patterns of Cu0.5Zn0.5AlxFe2-xO4

Fig.4.1 shows X-ray diffraction patterns of the samples Cu

ferrite samples (where x = 0.0 to 0.5)

0.5Zn0.5AlxFe2-xO4

(for x = 0, 0.1, 0.2, 0.3, 0.4, 0.5). As can be seen, all the samples can be indexed as

having a single phase cubic spinel structure. No impurity peak was noticed. The

breadth of the characteristic ferrite peaks is an indication of lower crystallite size of

the samples. The crystallite size have been estimated from the X-ray peak broadening

of (311) diffraction peak using the Scherrer formula [5]. For all the samples, the

crystallite size remained in the range of 25-30 nm. The value of the lattice constant ‘a’

of the cubic spinel calculated using the ‘CELL’ software have been listed in Table 4.1.

A decrease in lattice constant was observed with increase of Al3+ concentration in

samples. The decrease in lattice constant is justifiably be expected and can be



51

attributed to the substitution of smaller Al3+ ion (0.51 Å) for large Fe3+ ions (0.64 Å)

in the system Cu0.5Zn0.5AlxFe2-xO4. The bulk density (ρb) was calculated from the

weight and dimensions of the sintered samples using the relation, ρb = m/V [6], where

m is the mass and V is the volume of the samples. As obvious from the table 4.1, the

value of the bulk density decreased from 4.59 g/cm3 to 3.96 g/cm3 as the Al3+

concentration was increased from x = 0.0 to 0.5 in the series. The decrease in bulk

density is due to the fact that with the increase of Al, the porosity in the samples has

increased consistently as can be seen from the SEM micrographs given in Fig. 4.2. X-

ray density (ρx) of the samples was calculated using the relation, ρx = 8M/Naa3 given

by Smit and Wijn [7], where M is the molecular weight of the samples, Na is the

Avogadro’s number and a is the lattice constant. The number ‘8’ is included in the

formula as there are eight molecules per unit cell in the cubic spinel ferrite structure.

The value of ρx decreased from 5.41 g/cm3 to 5.07 g/cm3 with the increase in Al3+

contents in the sample series as the decrease in mass overtakes the decrease in volume

of the unit cell. It is noted that X-ray density of each sample is greater than the

corresponding bulk density which is an evidence of the presence of pores in the

samples. The porosity was found to increase from 0.151 to 0.219 in the series which is

direct evidence that the substitution of Al3+ for Fe3+

leaves relatively more empty

spaces in the samples. This is due to ceramic method which gives us large and non-

uniform particle size, on compacting results in the formation of voids and

subsequently lowers the density and increased the porosity.

Parameter x = 0.0

x = 0.1 x = 0.2 x = 0.3 x = 0.4 x = 0.5

a (Å)±0.001 8.385 8.383 8.340 8.317 8.266 8.211 V (Å3 589.53 ) 589.11 580.09 575.31 564.79 553.59 ρs (g/cm3 4.59 ) 4.38 4.27 4.15 4.10 3.96 ρx (g/cm3 5.41 ) 5.28 5.23 5.16 5.10 5.071 P (fraction) 0.151 0.170 0.183 0.195 0.196 0.247 Ms 75 (emu/cc) 72 68 56 50 38

E (eV) 0.450 0.452 0.393 0.437 0.462 0.440 Table 4.1 Lattice constant (a), lattice volume (V), sintered density (ρs), X-ray density

(ρx), porosity (P), saturation magnetization (Ms) and activation energy ( E) of

Cu0.5Zn0.5Fe2-xAlxO2 ferrite system



52

(a) (b)

(c) (d)

(e) (f)

Fig. 4.2 SEM micrographs of Cu0.5Zn0.5AlxFe2-xO4

Figure 4.2 (a-f) illustrates the representative micrographs of the Cu

with (a) x = 0.0, (b) x = 0.1, (c) x

=

0.2, (d) x = 0.3, (e) x = 0.4 and (f) x = 0.5.

0.5Zn0.5AlxFe2-

xO4 system that reveal surface morphology of the samples obtained using scanning

electron microscope. The images show that the grain size increases with increasing



53

Al3+ concentration and lies in the range of about 2-6 μm. The increased grain size in

the series refers to the more porous samples as is evident from the increased value of

porosity discussed earlier.

The magnetic hysteresis loops for the series of samples were obtained using

vibrating sample magnetometer. The results revealed that the value of saturation

magnetization decreased with the increase of Al3+

concentration as shown in the

Figure 4.3. We understand the trend as the substitution of a non-magnetic element

(Al) for a magnetic element (Fe) at the B-site of the cubic spinel structure has caused

the magnetization to decrease gradually [8].

Fig.4.3 Saturation magnetization plotted against Al3+

Figure 4.4 shows the temperature dependent variation in DC electrical resistivity

measured by two-probe method. The DC electrical resistivity increases as the Al

concentration.

3+

concentration increases for all the samples. This trend can be understood considering

the conduction mechanism in ferrites which takes place mainly through the hopping

of electrons between Fe2+ and Fe3+ at B-sites as explained by Verwey [9]. The

hopping probability depends upon the separation of ions involved and the activation

energy. As the distance between two metals ions at B-sites is smaller than the distance

between two metal ions, one at A-site and another at B-site, therefore the electron

hopping between A and B sites has a less probability as compared to hopping between

B-B sites. Hopping between A and B sites does not limit for the simple reason that

there are only Fe3+ ions at A site and only Fe2+

0.0 0.2 0.4 0.6 0.8 1.0

40

50

60

70

80

Ms (

emu/

g)

Al-Concentration (x)

Saturation Magnetization, Ms

ions preferentially occupy B site



54

during processing. Therefore, the deficiency of Fe2+ ions with increasing Al3+

concentration gives further reason for the increase of DC electrical resistivity. The

measured values of DC electrical resistivity at 293 K were found to vary from 2.16 x

106 Ωcm to 1.17 x 108 Ωcm as the concentration of Al3+ was increased from x = 0 to

0.5. High values of DC electrical resistivity and relatively easy preparation method

make ferrites an appropriate choice for the cores of intermediate and high frequency

electromagnetic absorbers.

The slopes of the linear plots of DC electrical resistivity as shown in Figure 4.4

determine the activation energy in the measured temperature range. In

Cu0.5Zn0.5AlxFe2-xO4

system, the values of activation energy were found to vary

between 0.393 to 0.462 eV. In ferrites, the activation energy is often associated with

the variation of mobility of charge carriers rather than their concentration. This

activation energy plays an essential role in overcoming the electrical energy barrier

experienced by the electrons during hopping process, which in turn, contributes

towards conductivity.

Fig.4.4 DC electrical resistivity plotted against temperature.

Figure 4.5 shows the variation of dielectric constant (ε′) with rise of frequency up

to 1MHz. The value of ε′ is higher at lower frequencies and is found to decrease with

increase in frequency. At high frequencies, particularly for the composition having x

= 0.3 to 0.5, the value becomes small, constant and independent of frequency [10].

The variation in dielectric constant is directly related with space charge polarization.

2.2 2.4 2.6 2.8 3.0 3.2 3.4 3.610

12

14

16

18

20

22

24 0.0 0.1 0.2 0.3 0.4 0.5

ln(ρ)

(Ω−c

m)

1000/T(K-1)



55

The presence of higher conductivity phases (grains) in the insulating matrix (grain

boundaries) of a dielectric produces localized accumulation of charge under the

influence of an electric field, results in space charge polarization [11]. A finite time is

needed for the space charge carriers to line up their axes parallel to an alternating

electric field. A continuous increase in field reversal frequency results in a point

where space charge carriers cannot remain preserved with the field and the alternation

of their direction lags behind the field, resulting in a reduction of dielectric constant of

the material [12]. In addition, space charge polarization also results from

inhomogeneous dielectric structure of the material as proposed by Maxwell and

Wagner in the form of two-layer model [13,14]. According to this model, space

charge polarization originates from large well conducting grains separated by thin

poorly conducting intermediate grain boundaries. In ferrites, polarization can also be

regarded as a similar process to that of conduction [15]. The hopping of electron

between Fe3+ and Fe2+

ions, results in the local displacement of electrons in the

direction of applied field that contributes towards polarization. When the frequency is

increased, polarization decreases until attaining a constant value. Beyond this critical

value of frequency, the electron exchange between the two cations, cannot follow the

alternating field.

Fig. 4.5 Dielectric constant plotted against frequency.

Predominance of species like Fe2+

4.5 5.0 5.5 6.0 6.5 7.0

200

400

600

800

1000

1200

Diel

ectri

c co

nsta

nt (ε

/ )

ln(f)

0.0 0.1 0.2 0.3 0.4 0.5

ions, oxygen vacancies, grain boundary defects

and voids significantly contribute to increase the dielectric constant at lower

frequencies [16]. At higher frequencies, any species contributing to polarizability lags

behind the applied field and hence the decreasing trend in dielectric constant is

witnessed.



56

The tangent of dielectric loss angle (tan δ) decreased with the increase of

frequency as shown in the Fig. 4.6. It is essential to note that the value of tan δ depend

on different factors like stoichiometry, Fe2+

content and structural homogeneity.

These factors, in turn, depend on the composition of the samples and their sintering

temperature [17]. The decrease of tan δ with an increase in frequency could be

explained on the basis of Koops phenomenological model [18].

Fig. 4.6 Tangent of dielectric loss angle plotted against frequency.

The tangent of dielectric loss angle (tan δ) decreased with the increase of

frequency as shown in the Fig. 4.6. It is essential to note that the value of tan δ depend

on different factors like stoichiometry, Fe2+

content and structural homogeneity.

These factors, in turn, depend on the composition of the samples and their sintering

temperature [17]. The decrease of tan δ with an increase in frequency could be

explained on the basis of Koops phenomenological model [18].

Fig. 4.7 Dielectric loss factor plotted against frequency.

4.5 5.0 5.5 6.0 6.5 7.00

5

10

15

20

25

30

35

Tan(δ

)

ln(f)

0.0 0.1 0.2 0.3 0.4 0.5

4.5 5.0 5.5 6.0 6.5 7.00

1000

2000

3000

4000

Diel

ectri

c lo

ss (ε

// )

ln(f)

0.0 0.1 0.2 0.3 0.4 0.5



57

An essential part of the total core loss in ferrites is termed as dielectric loss factor

(ε) [19]. Figure 4.7 shows the plot of frequency dependent dielectric loss factor. As

the number of hopping electrons increase, the extent of local displacement in the

direction of electric field increases, causing an increase in electric polarization, which

in turn enhances dielectric loss. The dielectric losses in ferrites are exhibited during

conductivity measurements, as highly conducting materials show high losses [20].

Therefore, the present ferrite series with relatively low losses might be useful in

technological applications at higher frequencies.

4.4 Conclusions Aluminum substituted CuZn-Ferrite materials prepared by conventional solid state

reaction technique exhibited single phase cubic spinel structure having nano-sized

crystallite size. The crystal lattice constant declines gradually from 8.385 Å to 8.211

Å, with the increasing Al3+ contents. This trend is attributed to the smaller ionic radius

of Al3+ as compared to Fe3+. The decrease in dc electric resistivity of the all the

samples with increasing temperature depicts the semiconductor like behavior of the

samples. The reason for decrease in saturation magnetization with increasing Al3+

contents in the CuZn-ferrite series could be understood considering the non-magnetic

nature of aluminum. The dielectric constant, tangent of dielectric loss and dielectric

loss factor, all showed decreasing trend with increasing frequency ensuring high

frequency applications of the Al3+ substituted CuZn-ferrite samples.



58

References 1- M. K. Shobana, S. Sankar and V. Rajendran, Mater. Chem. Phys. 113 (2009)

10.

2- I.H. Gul and A. Maqsood, J. Alloys Compd. 465 (2008) 227.

3- M. Ajmal and Asghari Maqsood, J. Alloys Compd. 460 (2008) 54.

4- A. Gonchar, V. Andreev, L. Letyuk, A. Shishkanov, and V. Maiorov, J. Magn.

Magn. Mater. 254/255 (2003) 544.

5- B. Zhou, Y. W. Zhang, C. S. Liao, F. X. Cheng, C. H. Yan, J. Magn. Magn.

Mater. 24 (2002) 70.

6- I. H. Gul, F. Amin, A. Z. Abbassi, M. Anis-ur-Rehman, A. Maqsood, J. Magn.

Magn. Mater. 311 (2007) 497.

7- J. Smit, H. P. J. Wijn, Ferrites, John Wiley, New York, 1959.

8- A. A. Sattar, J. Mater. Sci. 39 (2004) 451.

9- E. J. W. Vervey, J. H. De Boer, Rec. Trans. Chim, de Pays-Bas 55 (1936) 531.

10- R. Laishram, S. Phanjoubam, H. N. K. Sarma, C. Prakash, J. Phys. D: Appl.

Phys. 32 (1999) 2151.

11- M. Chanda, Science of Engineering Materials, vol. 3, The Machmillan

Company of India Ltd., New Delhi, 1980.

12- A. M. Shaikh, S. S. Bellad, B. K. Chougule, J. Magn. Magn. Mater. 195

(1999) 384.

13- J. C. Maxwell, Electricity and Magnetism, vol. 1 Oxford University Press,

Oxford, 1929 (Section 328).

14- K. W. Wagner, Ann. Phys. 40 (1913) 817.

15- I. T. Rabinkin, Z. I. Novikova, Ferrites, Izv Acad. Nauk USSR Minsk, 1960.

16- J. C. Maxwell, Electricity and Magnetism, vol. 2 Oxford University Press,

New York, 1973.

17- A. Verma, T. C. Goel, R. G. Mendiratta, P. Kishan, J. Magn. Magn. Mater.

208 (2000) 13.

18- C. G. Koops, Phys. Rev. 83 (1951) 121.

19- J. Zhu, K. J. Tseng, C. F. Foo, IEEE. Trans. Magn. 36 (2000) 3408.

20- A. S. Hudson, Marconi Rev. 37 (1968) 43.

Chapter 5 Fabrication and Characterization of Naostructured Magnetic Materials

59

Fabrication and Characterization of Nanostructured

Magnetic Materials By Sol-gel Combustion 5.1 Influence of temperature on the structural and

magnetic properties of Co0.5Mn0.5Fe2O4

ferrites

5.1.1 Motivation

Most of magnetic materials consist of metal or metal oxides. Preparation of

these materials in bulk form is a simple task, however a bit more challenging aspect

relates to phase purity, crystal structure and morphology which are responsible for

better performance of these functional materials. But when we reduce the crystal

dimension to nanometer scale, a new degree of complexity occurs to their synthesis.

Grinding method is simple and can be used for metal oxides (because most metals

are malleable), and for those areas of application where particle morphology and

phase purity are unimportant. By chemical synthesis in solution methodology, one

can control the particle size and their distribution, and uniformity in shape. In

chemical synthesis, temperature and concentrations are not the only parameters

which control the rate of reaction, type of precursors and the mechanism of the

reaction also play a major role [1].

CoFe2O4 ferrites are considered one of the most promising and

technologically important materials for high density recording media due to their

high coercivity (Hc), moderate saturation magnetization (Ms) and excellent chemical

stability [2]. Metal substituted cobalt ferrites are suitable for magneto mechanical

strain sensors and activators applications [3]. Many researchers have studied cobalt

ferrites with Mn substitution and investigated their magnetic, magneto-optical and

magneto-mechanical properties [4-8]. In this present work, we have synthesized

Co0.5Mn0.5Fe2O4 ferrites using the sol-gel auto-combustion method, a novel method

based on the combination of chemical sol-gel and combustion processes [17]. The as-

burnt powder was calcined at different temperatures (500°C, 600°C, 700°C, 800°C

and 900°C). The main objective of this work to study the size effect on the structural


60

and magnetic properties of Co0.5Mn0.5Fe2O4

5.1.2 Experimental details

ferrites calcined at different

temperatures.

Analytical grade ferric nitrate Fe(NO3)2.9H2O, cobalt nitrate

Co(NO3)2.6H2O, manganese nitrate Mn(NO3)2.4H2O, citric acid C6H6O7.2H2O and

ammonia were used as starting materials. Ferric nitrate, cobalt nitrate and manganese

nitrate in the molar ratios 2:0.5:0.5 were dissolved in deionized water. The mixed

solution was neutralized to pH 7 by adding liquid ammonia. After this the neutralized

solution was evaporated to dryness by heating at 100°C on a hot plate with

continuous magnetic stirring. As water evaporated, the solution became viscous and

finally formed a highly viscous gel. Increasing the temperature to up to about 200°C

led to the ignition of the gel. The dried gel burnt in a self propagating combustion

reaction until all the gels were completely burnt out to form a voluminous and fluffy

powder with large surface area. Finally the as-burnt powders were calcined at

different temperatures (500 – 900°C) for one hour. Experimentally it is observed that

all the samples showed combustion behavior and burn out completely to form a loose

powder.

In order to characterize the as-burnt and calcined powder, x-ray diffraction

with CuKα radiation (1.5406 Å, D-MaxII-A X-ray diffractometer) was used to

confirm the Co0.5Mn0.5Fe2O4

5.1.3 Results and discussions

phases. Magnetic properties were carried out at room

temperature by vibrating sample magnetometer (Lakeshore 7404).

Fig. 5.1 shows the X-ray diffraction patterns of as-burnt and samples calcined

at different temperatures (500, 600, 700, 800 and 900°C). The patterns show that all

the samples show cubic spinel structure.


61

Fig.5.1 XRD patterns of Co0.5Mn0.5Fe2O4, as-burnt and calcined at 500,600, 700,

800 and 900°C.

As we increase the calcination temperature, the crystallinity, lattice

parameters and crystallite size increase as shown in the Figs. 5.1, 5.2 and 5.3

respectively. At high temperature, we obtained some phases of α-Fe2O3

K Cos

D λβ θ

=

. This may

be due to the phase stability of iron oxide. The crystallite size was estimated by

considering the most intense peak (311) and using the Scherrer formula, [18]

.

Here D is the estimated crystallite size, β is the FWHM (full width half maximum)

and K is a constant. Shobana et.al [7] has also reported the same composition but

they showed amorphous behavior of as-burnt powder, instead of crystalline. In our

case both as-burnt and calcined powders have crystalline behavior. Fig. 5.3 shows the

relation between estimated crystallite size and calcination temperature. The size of

the crystallite size is observed to be increasing linearly with calcination temperature.

Calcination generally removes the lattice defects and strain but sometimes bind the

crystallites in clusters and increase their size [19].


62

Fig.5.2 Variation of lattice parameters with calcination temperatures.

Fig.5.3 Variation of crystallite size with calcination temperature

Fig. 5.4 shows the room temperature magnetic properties of as-burnt and

calcined samples. All the samples show ferromagnetic behavior. Figs. 5.5 and

5.6 reveal that with the increase in the calcination temperature, crystallite size

increases whereas the coercivity first increases and then decreases but saturation

magnetization remains within the range of 20 emu/g to 50 emu/g. Maximum

value of coercivity of 1470.25 Oe is obtained at 600 °C calcination temperature

with a crystallite size of 22.9 nm. Maaz et.al [20] has reported two reasons, first,

may be due to the expected crossover from single domain to multidomain

behavior with increasing size and secondly, may be from a combination of

surface anisotropy and thermal energies. The first effect is expected only in

CoFe2O4

500 600 700 800 900

8.380

8.385

8.390

8.395

8.400

8.405

8.410

8.415

Lattic

e pa

ram

eter

s (Å)

Calcined temperature (°C)

having particle size close to 50 nm [21-22] which is higher than the critical

500 600 700 800 90021

22

23

24

25

26

27

28

Crys

tallite

Size

(nm)

Calcined Temperature (°C)


63

size of 22.9 nm that we observed. The initial increase of coercivity with increasing

crystallite size is due to the dominant role of surface anisotropy as compared to bulk

anisotropy. But the most dominant role would be due to surface effect for smaller

particles [23]. Decrease of coercivity at larger crystallite size may be due to the

development of domain walls in the nanoparticles already reported by Maaz et. al

[20].

-12000 -8000 -4000 0 4000 8000 12000

-50

0

50-100-50

050

100-20-10

0102030

-50

0

50-40-20

02040

-50

0

50

M s(em

u/g)

H(Oe)

As burnt

5000C

6000C

7000C

8000C

9000C

Fig.5.4 Room temperature magnetic properties of Co0.5Mn0.5Fe2O4

calcined at different temperatures


64

Fig.5.5 Variation of coercivity with calcination temperature.

Fig.5.6 Coercivity as a function of crystallite size.

5.1.4 Conclusions Co0.5Mn0.5Fe2O4

500 600 700 800 900200

400

600

800

1000

1200

1400

1600H c

(Oe)

Calcined Temperature (°C)

nanoparticles were synthesized by sol-gel auto-combustion

method after adjusting metal nitrate to citric acid ratio to 0.5/1.0 and the pH to 7. The

structural and magnetic properties were studied as a function of calcination

temperature. The unit cell parameter a and crystallite size D increase linearly with the

21 22 23 24 25 26 27 28200

400

600

800

1000

1200

1400

1600

H c(O

e)

Crystallite Size (nm)


65

increase of temperature. As-burnt and calcined powders show crystalline behavior. A

maximum coercivity of about 1470 Oe with crystallite size 22.9 nm was obtained at

600°C which is higher than the value reported earlier [8]. Furthermore, no significant

change occurs in saturation magnetization as a function of calcination temperatures.


66

References 1. Sergey P. Gubin, Magnetic Nanoparticles, Wiley-VCH (2009).

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O. Kim and C.G. Kim, J. Magn. Magn. Mater., 316

6. O.F. Caltun, G.S.N. Rao, K.H. Rao, B. Parvatheeswara Rao, C.G. Kim, C.-O.

Kim, I. Dumitru, N. Lupu and H. Chiriac, Sens. Lett.

(2007) e618.

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7. B. Zhou, Y.W. Zhang, C.S. Liao, F.X. Cheng, C.H. Yan, L.Y. Chen and S.Y.

Wang, Appl. Phys. Lett.

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79

8. M.K. Shobana, S. Sankar, V. Rajendran, Material Science Communication,

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10. K. Suzuki, T. Yamazaki, Jpn. J.Appl. Phys. 27 (1988) 361.

11. L. Bouet, P. Namikawa Tailhades, A. Rousset, J. Magn. Magn. Mater., 153

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12. J. Daval, B. Bechevet, J. Magn. Magn. Mater., 129 (1994) 98.

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14. K. Shimokawa, H. Dohnomal, T. Mukai, H. Yamada, H. Matsuda, M. Daimo,

J. Magn. Magn. Mater. 154 (1996) 271.

15. T. Suzuki, J. Appl. Phys. 69 (1991) 4756.

16. G. Srinivas, S.C. Shin, Appl. Phys. Lett., 69 (1996) 3086.

17. A. Mali, A. Ataie, Scripta Materialia 53 (2005) 1065.

18. B.D. Cullity, S.R. Stock, Elements of x-ray diffraction analysis, Pearson

Education International (2007).

19. T.P. Raming, A.J.A. Winnusbst, C.M. Van Kats, P. Philipse, J. Colloid

Interface Sci. 249 (2002) 346.


67

20. K. Maaz, A. Mumtaz, S.K. Hasanain, A. Ceylan, J. Magn. Magn. Mater. 308

(2007) 289.

21. W.W. Schuele, Y.D. Deet Screek, W.W. Kuhn, H. Lamprey, C. Scheer (Eds),

Ultrafine particles, Wiley, New York (1963) p-218.

22. A.E. Berkowitz, W.J. Schuele, J. Appl. 30 (1959) 1345; C.N. Chinnasamy, B.

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23. M.A. Ahmed, E.H.E.I. Khawas and M.Y. Harson, J. Mater. Chem. 74

(2000)

567.


68

5.2 Low temperature synthesis of nanocrystalline Mn-Cu-Zn

Ferrites via sol-gel combustion method 5.2.1 Motivation

Many reports have been focused on the synthesis of controlled and fine

magnetic nanoparticles due to their technological and fundamental scientific

importance [1-2]. The structural and magnetic properties of nanoparticles have been

found to depend upon the particle size, which depends totally on the methods of

synthesis [3]. High temperature ceramic method is required for the completion of

solid-state reaction between the constituent oxides or carbonates. The particles

obtained by this method are large and non-uniform in size. These non-uniform

particles, on compacting results in the formation of voids and subsequently lower the

density. In order to overcome these difficulties wet sol-gel method has been used for

the production of homogenous, fine and grained ferrites. It is possible to produce fine

powders having high homogeneity and large surface area with chemical methods.

Several techniques including solid state reaction, coprecipitation, microemulsion and

ball-milling [4-8] have been used to synthesize ferrites at micro and nano levels but

so-gel combustion is a novel, low cost, energy-efficient and simple method which

contains a combination of chemical sol-gel and combustion processes. Combustion

process is based on gelling, salts of desired metals and organic fuels which gives

voluminous and fluffy powder with large surface area [9]. Also combustion reaction

is self-propagating producing an adiabatic temperature in the range of 1500-3000 K

[10-11] which is sufficient for the required phase of ferrimagnetic materials within a

very short time.

Among the soft ferrites, many polycrystalline ferrites have been studied for

several years due to their commercial importance as magnetic materials because of

low eddy currents, operatable at high frequency and dielectric loss. Due to these

physical properties they can be used in telecommunication, audio and video, power

transformers, radio frequency coils, rod antennas and read-write heads for high speed

digital tape [12]. The main aim of the present study is to synthesize Mn0.5Cu0.5-

XZnXFe2O4 (x=0, 0.1, 0.2, 0.3, 0.4, 0.5) ferrites at low temperature via sol-gel auto-

combustion method and studied the effect of zinc concentration on Cu-site.


69

5.2.2 Experimental details

Analytical grade ferric nitrate Fe(NO3)2.9H2O, manganese nitrate

Mn(NO3)2.4H2O, copper nitrate Cu(NO3)2.3H2O, zinc nitrate Zn(NO3)2.3H2O,

citric acid C6H6O7.2H2O and ammonia were used as starting materials. In

Mn0.5Cu0.5-xZnxFe2O4 compositions ferric nitrate, copper nitrate, zinc nitrate and

manganese nitrate according to their molar ratio were dissolved in deionized water.

The mixed solution was neutralized to pH 7 by adding liquid ammonia. After this the

neutralized solution was evaporated to dryness by heating at 100°C on a hot plate

with continuous magnetic stirring. As water evaporated, the solution became viscous

and finally formed a highly viscous gel. Increasing the temperature up to about

300°C led to the ignition of the gel. The dried gel burnt in a self propagating

combustion reaction until all the gels were completely burnt out to form a

voluminous and fluffy powder with large surface area. Experimentally it is observed

that all the samples showed combustion behavior and burn out completely to form a

loose powder.

In order to characterize the as-burnt powder, x-ray diffractometer with CuKα

radiation (1.5406Å, D-MaxII-A X-ray diffractometer) was used to confirm the

Mn0.5Cu0.5-xZnxFe2O4 phases. Magnetic properties were carried out at room

temperature by vibrating sample magnetometer (VSM) (Lakeshore 7404).


70

5.2.3 Results and discussions

Fig. 5.7 XRD patterns of Mn0.5Cu0.5-xZnxFe2O4 ferrites

Figure 5.7 shows the x-ray diffraction (XRD) patterns of as-burnt powder of

Mn0.5Cu0.5-XZnXFe2O4

K Cos

D λβ θ

=

samples. All compositions are of a single-phase spinel

structure, implying that temperature produced during combustion is sufficient for the

reaction of constituent nitrates to form the spinel ferrites. No additional phase was

detected. The broad and highest peak of (311) plane indicates lower crystallite size of

the synthesized samples. The estimated crystallite size was determined by

considering the most intense peak of (311) by Scherrer formula [13]

Here D is the estimated crystallite size β is the FWHM (full width half maximum)

and K is constant. The estimated crystallite size remains within the range of 21.7-

27.2 nm. The lattice constant (a) was calculated using the formula:

2 2 2hkla d h k l= + +


71

value of ‘a’ for each composition was calculated and tabulated in the Table5.1. This

shows that lattice parameter increases with Zn2+ concentrations (x). The increase in

lattice parameter may be due to the substitution of larger ionic radii of Zn2+ (0.74 Å)

for smaller Cu2+ (0.73 Å) ions in the system Mn0.5Cu0.5-XZnXFe2O4

0.0 0.1 0.2 0.3 0.4 0.521

22

23

24

25

26

27

28

Crystallite SizeLattice parameter

Zn Concentration

Crys

talli

te S

ize (n

m)

8.405

8.410

8.415

8.420

8.425

8.430

8.435

8.440

8.445

Latti

ce P

aram

eter

(A°)

. An increase in

lattice parameter is expected because larger ions are replacing smaller one. The

crystallite size datum point for Zn concentration x=0.5 shown in Fig. 5.8 is the lowest

of all. It is very interesting and rather difficult to explain without carrying out further

preparatory investigations.

Fig. 5.8 Variation of crystallite size and lattice parameter with zinc concentration of

Mn0.5Cu0.5-xZnxFe2O4

-6000 -4000 -2000 0 2000 4000 6000

-100

1020

-20

0

20

-20

0

20-40-20

02040

-50

0

50

-50

0

50

Hc (Oe)

0.0

0.1

M s (emu

/g)

0.2

0.3

0.4

0.5

ferrites

Fig. 5.9 Room temperature hysteresis loops of Mn0.5Cu0.5-xZnxFe2O4 ferrites


72

Saturation magnetization (Ms) and coercivity (Hc) of Mn0.5Cu0.5-XZnXFe2O4

ferrites were measured from M-H loops taken on a VSM at room temperature

presented in Table 5.1 and M-H loops are shown in figure 5.9. The variation in

saturation magnetization (Ms) and coercivity (Hc) as a function of Zn concentration

are shown in figure 5.8. As already reported [14] Zn ions have strong preference for

A-site (tetrahedral) substituted for Cu ions having strong preference for B-site

(octahedral), then Fe3+ ions will start transferring from A to B-site resulting in an

increase of magnetization of B-site. The coercivity decreases as the crystallite size

increases, attaining a minimum value of 46.32 Oe as shown in table 5.1. This

decrease at larger crystallite size could be due to three reasons. First crossover of

single domain to multiphase domain, second combined effect of surface and surface

anisotropy [15], third migration of Fe3+ ions from A to B-site. Among these three

processes, the dominant role would be due to the migration of Fe3+ ions from A to B-

site. Alex Goldman [16] predicted these materials as square loop ferrites but in our

case the squareness of the loops (Mr/Ms) decreases as we increase the

Zn2+ concentrations. Actually square-loop ferrites are of two types, one is

spontaneously square and second which becomes square after magnetic annealing

[17]. We synthesized the samples with sol-gel combustion without any further heat

treatment. It may be concluded that temperature during combustion is not sufficient

for spontaneous square behavior of Mn0.5Cu0.5-XZnXFe2O4

0.0 0.1 0.2 0.3 0.4 0.510

15

20

25

30

35

40

45

50 Saturation Magnetization Cercivity

H c (Oe)

M s (em

u/g)

Zn Concentration

0

50

100

150

200

250

300

350

400

450

nanoparticles.

Fig. 5.10 Variation of saturation magnetization (Ms) and coercivity (Hc) as a

function of zinc concentration


73

Concentration

(x)

Crystallite

Size

(nm)

Lattice

constant (a)

(Å)

±0.001

Coercivity

H

Magnetization

Mc

(Oe) s

(emu/g)

0.0 22.9 8.408 397.82 14.46

0.1 25.6 8.419 223.43 17.95

0.2 26.0 8.425 147.13 21.09

0.3 26.7 8.430 101.81 36.34

0.4 27.3 8.433 46.32 44.19

0.5 21.8 8.438 54.49 48.52

Table: 5.1 Crystallite size, lattice constants, coercivity and magnetization of

Mn0.5Cu0.5-xZnxFe2O4

5.2.4 Conclusions ferrites

Low temperature single phase polycrystalline Mn0.5Cu0.5-XZnXFe2O4

nanoparticles were successfully synthesized using the sol-gel combustion method.

The estimated crystallite size was calculated from the most intense peak (311) using

the Scherrer formula. The crystallite size of all the samples increases up to x=0.4

(zinc concentration). The lattice parameter ‘a’ increases with the increase of zinc

concentration due to larger ionic radius of Zn2+ compared to Cu2+ ions. Saturation

magnetization (Ms) and coercivity (Hc) increases and decreases as a function of zinc

concentration respectively already discussed in the previous section. The temperature

during combustion is sufficient for single phase but not sufficient for its square-loop

behavior.


74

References 1- A.P. Alivisatos, Science, 271 (1996) 933-937.

2- K.J. Klabunde, “Nanoscale Materials In Chemistry”, Wiley-Interscience,

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Edition, (2007) p-

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16- A. Goldman, Modern Ferrite Technology 2nd

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Business Media, Inc. (2006) p-106.


75

5.3 Low temperature synthesis and magnetic properties of

Mn0.5Cu0.5-x NixFe2O4

5.3.1 Motivation

nanoparticles via sol-gel

combustion method

In the recent past, many studies have been focused on the synthesis of controlled

magnetic nanoparticles, because of their technological and fundamental scientific

importance [1-2]. Magnetic nanoparticles exhibit very interesting structural, electrical,

optical and magnetic properties as compared to their corresponding bulk materials [3-

5]. Nanoferrites have been synthesized and studied due to their amazing electrical

and magnetic properties. These materials have high electrical resistivity, low eddy

current and dielectric losses, and can be used in telecommunication and transformers

[6].

The structural and magnetic properties of ferrites strongly depend on the

stoichiometry and methods of preparation. Several techniques, including solid state

reaction method, co-precipitation, micro-emulsion and ball milling [6-10], have been

used to synthesize ferrites at micro and nano levels but sol-gel auto-combustion is a

unique, economical, energy-efficient and simple method, which contains a

combination of chemically processed sol-gel and combustion processes. These

processes are based on gelling the salts of desired metals and some organic fuels,

which give us voluminous and fluffy powder after burning with large surface area

[11]. To develop spinel phases, we need high temperature for a long time, which is

usually obtained from laser radiation, a resistive heating coil and an electric arc, but

sol-gel combustion is a self-propagating reaction, producing an adiabatic temperature

in the range of 1500-3000 K [12], which is sufficient for the synthesis of

ferrimagnetic materials within a very short period of time.

Many researchers have prepared Mn1-xCuxFe2O4 ferrites using various methods

in order to corporate them for important technological applications [13-14]. The

substitution of Ni in these spinel ferrites could help to reduce the crystallite size and

in addition, owing to its ferromagnetic characteristics, might help to decrease the

coercivity and increase the saturation magnetization [15], in order to make these

ferrites appropriate for magneto-optical applications [16]. In the present work, we

have investigated the effect of Ni substitution (at Cu site), on the structural and


76

magnetic properties, considering the composition Cu0.5-XNiXFe2O4

5.3.2 Experimental

(x = 0, 0.1, 0.2,

0.3, 0.4 and 0.5) prepared by the sol-gel auto-combustion method. The work also

aims at investigating the spinel phase formation during a self propagating combustion

process.

Analytical grade ferric nitrate [Fe(NO3)2.9H2O], manganese nitrate

[Mn(NO3)2.4H2O], copper nitrate [Cu(NO3)2.H2O], nickel nitrate [Ni(NO3)2.H2O],

citric acid [C6H8O7] and ammonia (NH3) were used as starting materials, to prepare

the composition, Mn0.5Cu0.5-XNiXFe2O4 (x = 0, 0.1, 0.2, 0.3, 0.4 and 0.5). For this

purpose, nitrates of iron, copper, nickel and manganese, according to their

stoichiometric ratios were dissolved in de-ionized water. The mixed solution was

neutralized to pH 7 by adding proper amount of liquid ammonia (NH3). After that,

the neutralized solution was evaporated to dryness by heating at 100 °C on a hot plate

with continuous magnetic stirring. As water evaporated, the solution became viscous

and finally formed a highly viscous gel. Increasing the temperature up to about

300 °C led to the ignition of the gel. The dried gel burnt in a self propagating

combustion reaction until all the gel was completely burnt out to form a voluminous

and fluffy powder with large surface area. Experimentally, it was observed that all

the samples showed combustion behavior and burnt out completely to form a loose

powder.

In order to characterize the as-burnt powder, X-ray diffractometer (XRD) with

CuKα radiation (1.5406 Å, D-MaxII-A X-ray diffractometer), was used to confirm

the Mn0.5Cu0.5-XNiXFe2O4 phases. Magnetic properties were determined at room

temperature using a Lakeshore-7404, vibrating sample magnetometer (VSM).


77

Fig. 5.11 X-ray diffraction patterns of as-burnt Mn0.5Cu0.5-XNiXFe2O4

5.3.3 Results and discussion

powders

Figure 5.11 shows the as burnt XRD patterns of Mn0.5Cu0.5-XNiXFe2O4

The lattice constant of cubic Mn

samples

prepared by varying Ni concentrations from x = 0.0 to 0.5 with a step increment of

0.1. All the compositions revealed single phase spinel structure, implying that the

temperature produced during self burning was sufficient for the reaction of

constituents to form the desired spinel ferrites. However, the intensity of the

diffraction peaks was seemed to decrease as the Ni contents were increased. It could

be inferred that, although a sufficiently high temperature was produced during self

combustion process, yet time duration for which it persisted, was not sufficient, for

the Ni substituted spinel ferrite structure, to develop in a well-oriented manner.

0.5Cu0.5Fe2O4 (JCPDS No. 01-074-2072) is

8.410 Å while that of cubic NiFe2O4 is 8.258 Å (JCPDS No. 01-074-1913).

Therefore, one can expect that the lattice constant of Mn0.5Cu0.5-XNiXFe2O4 should

decrease with the increase in concentration of Ni from x = 0 to 0.5. However, the

XRD patterns of our samples did not show any consistent trend in diffraction peaks

shifting to either lower or higher angles. Therefore, the lattice constant ‘a’ showed a

non-consistent trend, as shown in Fig. 5.12. To calculate exact lattice parameters,

high-angle x-ray diffraction (2θ ≥ 90°) is usually required for polycrystalline single

phase ceramics. In the present work, we made no attempt for high-angle diffraction,


78

so it was very difficult

to justify the non-consistent trend observed at low-angle x-ray diffraction.

The crystallite size was estimated as 126, 111, 98, 67, 57 and 51 nm for x = 0,

0.1, 0.2, 0.3, 0.4 and 0.5 of nickel concentration in the series of samples respectively,

as depicted by Fig. 5.12, evaluated by considering the most intense diffraction peak

(311), using the Scherrer formula [17]. The trend of an increase in the lattice

parameters on the increased substitution of smaller sized Ni2+ radii (0.69 Å) in place

of larger sized Cu3+

(0.73 Å), as depicted in Fig. 5.12, is contrary to general

expectations, and difficult to support without further investigations.

Fig. 5.12 Variation of lattice constant and crystallite size with Ni concentration

Figure 5.13 shows the magnetic hysteresis loops for all the samples obtained

using VSM with an in-plane applied field of ± 5 kOe. The saturation magnetization

and coercivity of the samples as a function of nickel concentration are shown in

Fig.5.14. The decrease of coercivity and increase in magnetization with nickel

concentration might very well be understood considering the soft ferromagnetic

nature of Ni2+ ions when replaced with the diamagnetic Cu2+ ions. Figure 5.15 shows

the variation of Hc and Ms with the increasing crystallite size. It is well known that

coercivity depends on many factors such as grain size, grain shape, crystal defects

and packing density [16], but the most important one is the crystallite size [18-19].

All the magnetic hysteresis loops shown in Fig.5.13 exhibit characteristic

ferromagnetic behavior which is an evidence that the grain size of Mn0.5Cu0.5-

XNiXFe2O4

0.0 0.1 0.2 0.3 0.4 0.58.40

8.42

8.44

8.46

Lattice ConstantCrystallite Size Linear Fit of B

Ni Concentration (x)

Latti

ce C

onst

ant (Å

)

40

60

80

100

120

Crys

tallit

e siz

e (n

m)

particles has not reached the superparamagnetic threshold, which is

expected when the crystallite size becomes less than a certain critical value [20].


79

Fig.5.13 RT hysteresis loops for Mn0.5Cu0.5-xNixFe2O4

ferrites with varying Ni

concentration

Fig. 5.14 Variation of coercivity and saturation magnetization as function of Ni

concentration

-6000 -3000 0 3000 6000-10

01020-300

30-30

030-40

040

-500

50-50

050

0.0

Hc(Oe)

0.1

M s(e

mu/

g)

0.2

0.3

0.4

0.5

0.0 0.1 0.2 0.3 0.4 0.5

100

200

300

400 Coercivity saturation Magnetization

Ni Concentartion (x)

Coer

civity

(Oe)

10

20

30

40

50

Ms

(em

u/g)


80

Fig. 5.15 Variation of coercivity and saturation magnetization as a function of

crystallite size

Concentration (X)

Crystallite size (nm)

Lattice constant

(Å) ±0.001

Coercivity Hc

Magnetization M (Oe) s (emu/g)

0.0 126 8.408 397.82 14.46

0.1 111 8.424 120.27 30.02

0.2 98 8.419 108.99 33.76

0.3 67 8.441 97.62 41.15

0.4 57 8.453 89.74 50.02

0.5 51 8.442 87.20 52.71

Table.5.2 Crystallite size, lattice constants, coercivity and magnetization of

Mn0.5Cu0.5-xNixFe2O4

5.3.4 Conclusions

ferrites

Nano-crystalline, Mn0.5Cu0.5-XNiXFe2O4 ferrites, varying x from 0 to 0.5 have

been synthesized successfully by sol-gel auto-combustion method. XRD revealed the

structure as single phase cubic spinel. The presence of Ni2+ ions did not show a

consistent trend in diffraction peaks shifting to either lower or higher angles. The

crystallite size was decreased as the Ni contents were increased in Mn0.5Cu0.5Fe2O4

50 60 70 80 90 100 110 120 130

100

150

200

250

300

350

400

Saturation magnetization Coercivity


Hc(

Oe)

10

20

30

40

50

Ms(

emu/

g)

.


81

The coercivity was decreased and saturation magnetization was increased in the

series, which was attributed to the substitution of ferromagnetic Ni2+ contents in

place of diamagnetic Cu2+ ions. Minimum value of coercivity (87.20 Oe) was

observed for the composition Mn0.5Ni0.5Fe2O4.


82

References 1- A.P. Alivisatos, Science, 271 (1996) 933.

2- K.J. Klabunde, “Nanoscale materials in chemistry”, Wiley-interscience, New

York, 2001.

3- S.A. Majetich and Y. Ying, Science, 284 (1999) 470.

4- C.B. Murray, C.R. Kagan and M.G.Bawendi, Science, 270 (1995) 1335.

5- A.J. Zarur and J.Y. Ying, Nature, 403 (2000) 65.

6- Z.X. Tang, C.M. Sorensen, K.J. Klabunde, G.C. Hadjipanayis, J. Colloid

Interface Sci. 146 (1991) 38.

7- J.A. Lopez Perez, M.A. Lopez Quintela, J. Mira, J. Rivas, S.W. Chales, J.

Phys. Chem. B, 101 (1997) 8045.

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Phys. Lett., 80 (2002) 1616.

9- M.U. Rana. Misbah-ul-Islam, T. Abbas, Solid State Commun. 126 (2003) 129.

10- N.S. Gajbhiye, G. Balaji, M. Ghafari, Phys. Status Solidi (a), 189 (2002) 357.

11- M. Mali, A. Ataie, Scripta Materialia, 53 (2005) 1065.

12- Kashinath C Patil, Singanahally T Aruna and Sambandan Ekambaram,

Current Opinion in Solid State and Material Science, 2 (1997) 158.

13- Jian-Jun Li, Wei Xu, Hong-Ming Yuan, Jie-Sheng Chen, Solid State

Commun.,131 (2004) 519.

14- M.U. Rana. Misbah-ul-Islam, T. Abbas, Solid State Commun. 126 (2003) 129.

15- Amarendra K. Singh, Abhishek K. Singh, T.C. Goel, R.G. Mendiratta, J.

Magn. Magn. Mater., 281 (2004) 276.

16- M.K. Shobana, S. Sankar and V. Rajendran, Mater. Chem. Phys., 113, (2009)

10.

17- B.D. Cullity, S.R. Stock, Elements of x-ray diffraction analysis, Pearson

Education International, 2007.

18- K. Maaz, A. Mumtaz, S.K. Husnain and A. Ceylan, J. Magn. Magn. Mater.,

308 (2007) 295.

19- In: W.W. Schude, Y.D. Deet, W.W. Screek, H. Kuhn, C. Lamprey and Sheer,

Editors, Ultrafine Particles, Wiley, New York, (1963) 218.

20- A.E. Bekowitz and W.J. Shuele, J. Appl. Phys. 30

(1959) 345.

Chapter 6 Fe3O4 thin films

83

Fe3O4

6.1 Effect of Temperature on Structural and Magnetic

Properties of Laser Ablated Iron Oxide Deposited on

Si(100)

thin films on Si(100) substrate with pulsed laser deposition technique

(This work is published in Chinese Physics Letter, 2009)

6.1.1 Motivation Fe3O4 is predicted as half-metal (i.e, majority spin electrons are metallic and

minority are semiconducting) and considered as promising material for

magnetoelectronic or spintronic (spin transport electronics or spin based electronics, it

is not the electron’s charge but the electron’s spin that carries information) with a very

high Curie temperature (860 K) and 100 % spin-polarization (i.e, if every mobile

electron in the contact material has the same electron spin orientation). It has low

electrical resistivity at room temperature (10-3Ω-cm) and the structure changes from

cubic to monoclinic at 120 K temperature, called as Verwey transition temperature.

Spin life in semiconductor materials is high as compared to metals but there is a

problem to create a spin polarized population of the charge carriers in semiconductors

by injection of spin polarized currents. This can be done in dilute magnetic

semiconductors but they have low Curie temperature as compared to half metallic

materials like Fe3O4 or CrO2

Previously, fabrication of Fe

etc. [1-3].

3O4 on Si(100) has been achieved by different

techniques [4-5]. Molecular Beam Epitaxy (MBE) is very well established for epitaxy

of oxide thin films including Fe3O4, but this technique is exuberantly more costly

than pulsed laser deposition (PLD). In this study, we have fabricated Fe3O4 thin films

on Si(100) substrates at different temperatures (from room temperature to 450°C) by

pulsed laser deposition, and investigated the effect of annealing and deposition

temperature on the structural and magnetic properties of Fe3O4 thin films. Phase

analysis of the films was carried out by X-ray diffraction with CuKα radiation. Grain

size, lattice strain and lattice constants were measured by the Williamson-Hall plot [6],

assuming that the peak shapes are Lorentzian. Crystal structure determination was

performed using the powderX software [7]. Film thickness and magnetic properties


84

were determined from scanning electron microscopy (SEM) and vibrating sample

mangetometery (VSM) respectively.

6.1.2 Preparation of Fe3O4

The films were deposited on Si(100) substrates by ablating a commercially purchased

Fe

thin films

3O4 target having one inch diameter and 5 mm thickness. Before deposition, the

Si(100) sustrates were subjected to chemical cleaning using acetone and isopropanol

in an ultrasound bath. The cleaned substrates were then preheated in the deposition

chamber at 500°C for 30 minutes under a vacuum of 10-7 torr remove oxides.

The deposition procedure lasted for 20 minutes. For different expeiments, the

substrate was kept at different temperature from room temperature to of 450°C. In

each case, a background pressure of 10-6 torr was achieved after adjusting the oxygen

flow rate to 0.2 sccm. The pulse repetition rate, energy density of the Nd:YAG laser

(Ekspla Nl-303) and target to substrate distance were set at 10 Hz, 1.3 J/cm2 and 35

mm respectively. The laser wavelength was 266 nm.

Subsequent to deposition, film thickness was determined using scanning

electron microscopy (SEM) (Hitachi S-4800). Furthermore, the room temperature

(R.T) deposited films were annealed at 350, 400 and 450°C for 1 hr. in base pressure

of 1x 10-6 torr. The crystalline structure and phases were determined by X-ray

diffraction (XRD) (Rigaku D/Max-Rc MPA) using CuKα radiation. Finally, the room

temperature magnetic properties were measured by vibrating sample magnetometry

(VSM) (Ricken Denshi, Japan).

The depositing and annealing conditions of Fe3O4 thin films on Si(100) substrates are

summarized as follows:


85

Target Materials Substrates Depositing

temperature and

time

Annealing

temperature/

time

Fe3O Si(100) 4 Room temperature

for 20 minutes

300°C

400°C

450°C

for 1hr.

Fe3O Si(100) 4 350°C

400°C

450°C

for 1hr.

X

Table 6.1 Target materials with deposition and annealing temperature and time

6.1.3 Characterizations The iron oxide thin films prepared by pulsed laser deposition (PLD) were

studied by different characterization techniques in order to attain their structural and

magnetic properties. X-ray diffractometry (XRD) and vibrating sample magnetometry

(VSM) were used to determine the structural and magnetic properties of the material

respectively. Thickness calibration of the deposited and annealed thin films were

performed by atomic force microscopy and subsequently confirmed by cross-sectional

images obtained using scanning electron microscopy (SEM).

6.1.4 SEM for thin-film thickness determination

Fig. 6.1 shows the cross-sectional images obtained using scanning electron

microscopy (SEM) of annealed at 450°C and as-deposited films at 450°C as follows:

Fig. 6.1 SEM images of (a) Fe3O4 thin film annealed at 450°C and (b) as-deposited

film at 450°C.


86

It is clear from the images that annealed and as-deposited films have 117 and 143nm

thickness.

6.1.5 X-ray diffraction (XRD) analysis

Fig. 6.2 and 6.3 show the θ2θ scan XRD patterns of Fe3O4 thin films from

room temperature to 450°C on Si(100) substrates. The lattice parameters were

calculated using powderX software [7]. The grain size and lattice strain were

calculated from the Williamson-Hall plot using the following equation [8]:

B cos θ = kλ/L+S sin θ

where B=Full width half maximum, L= volume averaged grain size, k is the Scherrer

constant, in our case taken to be 1, S = lattice strain and λ = wavelength of CuKα

radiation.

As the annealing temperature was raised from 300°C to 450°C, the diffractograms

provide evidence of 99% single phase polycrystalline Fe3O4 thin films, except for Fe

peak. The crystallanity of the films increased as we increased the annealing

temperature and the peaks became sharper, showing, as expected, that the grain

increases with increasing annealing temperature. The peak positions also shifts to

higher angles as compared to bulk Fe3O4. This shifting is related to uniform lattice

strains in the plane of the films [8] compared to bulk sample. The situation was, in

fact, different for different deposition temperatures. The as-deposited films showed

smaller lattice strains as compared to annealed films. However, the trends in the unit

cell and crystallite size are identical in the annealed and as-deposited films. The

results are presented in Table 1 and 2. The increasing trend for the crystallite size has

also been reported by Tangel et al [9] and Parames et al [10].


87

(311)(111) (440)

R.T

(511) (440)(400)(222)

(220)(111)

Fe

Si

Si

Si

300oC

Si

(511)(440)

(400)

(222)

(220)

(111)400oC

Fe

20 30 40 50 60 70 80 90

(311)

(311)

(440)(511)

(311)(220)

Fe450oC

2-Theta (degree)

Inte

nsity

(arb

.uni

t)

(111)

Fig.6.2 X-ray diffraction patterns of Fe3O4 thin films on Si(100) substrates

deposited at room temperature and annealed at the shown temperatures


88

Fe

350oC(111)

(440)

Fe(440)(111) 400oC

Si

(311)

20 30 40 50 60 70 80 90

(440)(311)(111) 450oCFe

Inte

nsity

(arb

.uni

t)

2-Theta(degree)

(222)

Fig. 6.3 XRD diffraction patterns of as deposited thin films


89

Annealing

temperature

(°C)

Lattice Constant a

(Å)

±0.001

Lattice Strain S Estimated

Crystallite Size

(nm)

300 8.375 0.012 26.5

400 8.369 0.017 77.0

450 8.361 0.027 184.9

Table 6.2 XRD and VSM analysis of annealed Fe3O4

Depositing

Temperature

(ºC)

thin films on Si(100) substrates.

Lattice Constant a

(Å)

±0.001

Lattice Strain S Estimated Crystallite

Size

(nm)

350 8.369 .0035 18

400 8.364 .0011 22

450 8.348 .0010 30

Table 6.3 XRD and VSM analysis of as-deposited Fe3O4

6.1.6 Magnetic Properties

thin films on Si(100)

substrates

Figures 6.4 and 6.5 show the R.T. magnetization hysteresis behavior of

annealed and as-deposited samples. The figures clearly show hysteretic behavior

suggesting their R.T. ferromagnetic property. The variations of coercivity Hc(Oe)

with grain size are shown in Figures 6.6 and 6.7. The saturation magnetization (Ms)

and coercivity (Hc) as a function of annealing and depositing temperature (for both

the annealed and as-deposited samples) are also summarized in Tables 1 and 2. In

both cases the coercivity decreased with increasing annealing and depositing

temperatures and with increasing volume averaged crystallite size. At an annealing

temperature of 450°C, we obtained high saturation magnetization (Ms) as compared

to the bulk magnetization of Fe3O4 single crystals (471emu/cc) [11]. Kennedy and


90

Stampe have reported a high saturation magnetization for the magnetite thin films

grown on Si(100) substrate [12]. However, these authors used Fe as the target

material and suggested that the increased saturation magnetization could be due to the

increased Fe content in the thin films, which due to its amorphous structure, did not

show up their XRD results. In our XRD patterns, sharp peaks for crystalline Fe

become visible in the region 2θ ~45°. The enhanced magnetization is very likely due

to the presence of iron-rich centers. The hysteretic loops, contrarily, indicate a single

phase magnetic material but it is possible that the coercivity of the Fe is much too

small (as compared to the M-H loop step size) to distinctly appear in M-H loop.

Fig.6.4 Inplane magnetization curves of annealed thin films with A, B and C

representing samples annealed at 300, 400 and 450°C respectively.


91

Fig.6.5 Inplane magnetization curves of as-deposited thin films with A, B and C

representing deposition temperatures of 350, 400 and 450°C respectively.


92

25 50 75 100 125 150 175 200

325

350

375

400

425

Hc-Coercivity


H c(Oe)

Fig.6.6 Variation of Hc

18 20 22 24 26 28 30

275

300

325

350

375

400

425

450

Hc-coercivity)


H c(Oe)

with crystallite size of annealed thin films

Fig.6.7 Variation of Hc with crystallite size of as-deposited thin films.


93

6.1.7 Conclusions

We have fabricated Fe3O4 thin films by PLD at different temperatures (from room to

450°C) on Si(100) substrates and studied the effect of annealing and depositing

temperatures on the structural and magnetic properties of Fe3O4 thin films. XRD

patterns of both series showed the cubic inverse-spinel structure with different

orientations.

Annealing increased the crystallinity of the samples. As far as the magnetic

properties are concerned, we obtained ferromagnetic behavior of all the thin films but

obtained a surprisingly high magnetization of 854 emu/cc at 450°C (annealed

temperature) which is higher than the bulk value (471 emu/cc) of Fe3O4

. This may be

due to iron rich regions within the films as already reported. By increasing the

annealing and depositing temperatures, the lattice parameters and coercivity decreased,

while the volume average crystallite size increased.


94

References

1- F.C. Voogt, T.T.M. Palstra, L. Niesen, O.C. Rogojanu, M.A. James and T.

Hibma: Phys. Rev. B 57 (1998) R8107-R8110.

2- M. Ferhat and K. Yoh: Appl. Phys. Lett.90(2007) 112501-1-3

3- V. Dediu, E. Arisi, I. Bergenti, A. Riminucci, M. Solzi, C. Pernechele and M.

Natali: J. Magn. Magn. Mater. 316 (2007) e721-e723

4- D.T. Margulies, F.T. Parker, F.E. Spada, R.S Li J Goldman, R. Sinclair and

A.E Berkowitz, Phy. Rev. B 53 (1996) 9175.

5- W.F.J. Fontijn, R.M. Wolf, R. Metselaar, P.J. vander Zaag, Thin Solid Fims,

292 (1997) 270.

6- G.K. Williamson, W.H. Hall, Acta Mettall. 1 (1953) 22.

7- Cheng Dong, PowderX, J. Appl. Cryst. 32 (1999) 838.

8- B.D. Cullity, S.R Stock, Elements of X-ray diffraction 3rd

9- J. Tang, K.Y. Wang, W.J. Zhou, Appl. Phy. 89 (2001) 7690.

Edition, Prentice

Hall (2002).

10- M.L. Parames, J. Mariano, M.S. Rogalski, N. Popovici, O. Conde, Material

Science and Engineering B, 118 (2005) 246.

11- T. Hibma, F.C. Voogt, L. Niesen, P.A.A. Van der heijden, W.J.M. de Jonge,

J.J.T.M. Donkers, P.J.V. Zaag, J. Appl. Phys. 85 (1999) 5291.

12- R.J Kennedy and P.A. Stampe, J. Phys. D, Appl. Phys. 32 (1999) 16.


95

6.2 Effect of annealing time on structural and magnetic

properties of laser ablated oriented Fe3O4

6.2.1 Motivation

thin films

deposited on Si(100)

Fe3O4 is promising and technological important material due to half-metallicity.

Numerous studies have been reported of depositing highly oriented expitaxial or

polycrystalline Fe3O4 thin film techniques such as molecular beam epitaxy (MBE) [1-

2], electron beam ablation [3], reactive sputtering [4-6] and pulsed laser deposition [7-

8]. It is well known that in addition to the desired Fe3O4, several other phases can co-

exist such as Fe2O3, FeO and Fe according to the specific deposition condition [9-11].

However it is still difficult to grow them with well-defined composition with pulsed

laser deposition from different targets.

In the present study, we deposited Fe3O4 films on Si(100) substrates at 450°C by

pulsed laser deposition (PLD) technique and systematically investigated the effect of

annealing time on the structural and magnetic properties of Fe3O4

6.2.2 Preparation

thin films.

Fe3O4 thin films were grown on Si(100) substrates by pulsed laser deposition

(PLD) from commercially purchased Fe3O4 target. Before deposition the substrates

were cleaned with isopropanol in an ultrasonic bath for 20 minutes and then annealed

at 500°C for 30 minutes under a vacuum of 10-7 torr. A Nd:YAG laser (EKSPLA) of

wavelength 248 nm and pulse duration 3-6 nm was used to ablate the target. The

target was rotated at the rate of 10 rpm to avoid any crack formation on the target. The

pulse repetition rate was adjusted at 10 Hz and the energy density of the laser beam at

the target was 1.3 J/cm2. Deposition was carried out at a substrate temperature of

450°C for 20 minutes under working pressure of 10-6 torr after adjusting the flow rate

of oxygen to 0.3 sccm while the target to substrate distance was held fixed at 36 mm.

We annealed the films for 30, 60 and 90 minutes at the same temperature (450°C) and

pressure of 10-7 torr without oxygen flow. After deposition and annealing, the

substrates were cooled at the rate of 5°C/min. The film thickness, crystal structure and

magnetic properties were determined by scanning electron microscopy (SEM), X-ray

diffractometry (XRD) and vibrating sample magnetometry (VSM) respectively.


96

6.2.3 Results and discussion Fig.6.8 shows the cross-sectional view of film thickness confirms 143 nm. In

order to clarify the additional phases of iron oxide in Fe3O4 thin films with annealing

time we performed x-ray diffraction (XRD) measurement of these films. Fig. 6.9

shows the XRD patterns of 143 nm Fe3O4

Fig.6.8 Film thickess measured by scanning electron microscopy (SEM)

thick films before and after annealing. In

every pattern, we obtained Si(100) peak originating at 2θ ~ 69.4° (not shown in the

diagram due to some experimental problems). It is clear that all the films are grown

with preferred orientation in the [111] direction with a cubic structure. The lattice

parameter as a function of annealing time show that with increasing annealing time

the lattice parameter deviated from bulk material (8.396 Å) and are presented in Table

6.4, this may be due to the substrate induced strain in the film.

The size of the crystallite with annealing time increases due to the increase in

surface mobility [12], this is well known effect. The most interesting point is that at

90 minutes annealing we got pure preferential growth of the film in [111] direction.

As reported by Shailji Tiwari et al. [13], this may be due to the large lattice mismatch

between the films and substrates. As the substrate control over the film growth is

weak, the preferred orientation is determined by the thermodynamically stable state

having a minimum internal energy [14].


97

Fig. 6.9 XRD-patterns of Fe3O4

Annealing

time

(minutes)

thin films at different annealing time

Lattice

constants

(Å)

±0.001

Grain

Size

(nm)

Lattice

Strain

Saturation

Magnetization

(emu/cc)

Coercivity

Hc

(Oe)

0.0 8.378 18 .0084 330.34 303

30 8.377 120 .0075 366.30 306

60 8.367 127 .0068 276.14 272

90 8.366 157 .0010 335.00 315

Table 6.4 Lattice parameters, crystallite size, lattice strain, saturation

magnetization and coercivity of all samples

Fig.6.10 shows the room temperature magnetization hysteresis behavior for all the

films. We obtained a low saturation magnetization in all the films as compared to bulk

material. This lowering of saturation magnetization may be due to the presence of

antiphase boundaries between the films and substrates. Actually antiphase boundaries

even in epitaxial Fe3O4 films come from the nucleation of islands when the films are

deposited on substrates. Voogt et.al [1] showed that antiphase boundaries are formed


98

in the first monolayer, with a fixed domain size as subsequent layers are deposited but

Eerenstein et. al [15] have reported that domain size depends on the thickness of the

film. They reported that the domain size increases significantly with film thickess, and

therefore with deposition time. The increase in domain size with thickness has two

possibilities. One is that small domains are formed in the first monolayer and larger

domains grow on top of these as film thickness increases. Secondly the antiphase

boundaries migrate laterally during the growth process. Similar results have also been

reported earlier by Tiwari et. al [13].

-6000 -4000 -2000 0 2000 4000 6000-400

-300

-200

-100

0

100

200

300

400 zero minute

H (Oe)

M (e

mu/

cc)

-6000 -4000 -2000 0 2000 4000 6000

-400

-300

-200

-100

0

100

200

300

400

30 minutes

H (Oe)

M (e

mu/

cc)

-6000 -4000 -2000 0 2000 4000 6000

-300

-200

-100

0

100

200

300

60 minutes

H(Oe)

M (e

mu/

cc)

-6000 -4000 -2000 0 2000 4000 6000-400

-300

-200

-100

0

100

200

300

400

90 minutes

H (Oe)

M (e

mu/

cc)

Fig. 6.10 Vibrating sample magnetometry(VSM) of Fe3O4

6.2.4 Conclusions

thin films annealed at 0, 30, 60 and 90 minutes.

In conclusion, Fe3O4 thin films were deposited with pulsed laser deposition

technique on Si(100) substrates at 450°C for 30, 60 and 90 minutes annealing time.

The XRD patterns of the films imply the single phase spinel cubic structure with

[111] orientation at 90 minutes annealing. It was found that the grain size and lattice

strain increased and decreased with annealing time respectively. Magnetization results

showed ferromagnetic behavior for all the films, with saturation magnetization lower


99

than the bulk material may due to the presence of antiphase boundaries between films

and substrates as already explained in the previous section.


100

References 1- F.C. Voogt, T.T.M. Palstra, L. Niesen, O.C. Rogojanu, M.A. James and T.

Hibma: Phys. Rev. B 57 (1998) R8107-R8110.

2- M. Ferhat and K. Yoh: Appl. Phys. Lett.90(2007) 112501-1-3

3- V. Dediu, E. Arisi, I. Bergenti, A. Riminucci, M. Solzi, C. Pernechele and M.

Natali: J. Magn. Magn. Mater. 316 (2007) e721-e723

4- D.T. Magulies, F. T. Parker, F.E. Spada, R.S. Goldman, J. Li, R. Sinclair and

A. E. Berkowitz: Phys. Rev. B 53 (1996) 9175-9187.

5- C. Park, Y. Shi, Y. Peng, K. Barmak, J.-G. Zhu, D.E. Laughlin and R.M.

White: IEEE Trans. Magn.39 (2003) 2806-2808

6- H. Liu, E.Y. Jiang, H. L. Bai, R. K. Zheng, H. L. Wei and X.X. Zhang: Appl.

Phys. Lett. 83 (2003) 3531-3533.

7- G.Z. Gong, A. Gupta, G. Xiao, W.Qian, V.P. Draivid: Phys. Rev. B 56 (1997)

5096-5099.

8- M.L. Parames, J. Mariano, Z. Viskadourakis, N. Popovoco. M.S. Rogalski, J.

Giapintzakis and O. Conde: Appl. Surf. Sci. 252 (2006) 4610-4614.

9- C. Park, Y. Shi, Y. Peng, K. Barmak, J.-G. Zhu, D.E. Laughlin and R.M.

White: IEEE Trans. Magn.39 (2003) 2806-2808

10- E. Lochner, K.A. Shaw, R.C. Dibari, W. Portwine, P. Stoyonov, S.D. Berry

and D. M. Lind: IEEE Trans. Magn. 30

11- F.C. Voogt, T. Fujii, P.J.M. Smulders, L.Niesen, M.A. James and T. Hibma:

Phys. Rev. B 60 (1999) 11193-11206

12- Kiyotaka Wasa, Makato, Kitabatake, Hideaki Adachi, Thin Film Materials

Technology-Sputtering of compound materials, William Andrew Pub.-

Springer, 2004.

13- Tiwari S, Choudhary R J, Prakash R and Phase D M 2007 J. Phys.: Condens.

Matter 19 176002.

14- Tiwari S, Prakash R, Choudhary R J and Phase D M 2007 J. Phys. D 40 4943-

4947.

15- W. Eerenstein, T.T. Palstra, T. Hibma, Phy. Rev. B 68 (2003) 014428.

Conclusion

101

Conclusions

Owing to their diversity of compositions and properties, ferrites have always

been considered quite important materials, as far as their applications in electronic

and telecommunication industries are concerned. The work described in this thesis is

an experimental study, which is carried out to investigate the structural, electrical and

magnetic properties of some technologically important ferrite materials. The

preparation methods always play a key role in imparting desired properties to a

material. In this study, ferrite materials were prepared using conventional solid state

reaction method, state of the art sol-gel auto-combustion technique and pulsed laser

deposition (PLD). Al3+ doped Cu-Zn ferrites were prepared by ceramic method, Ni2+

and Zn2+ substituted Mn-Cu ferrites and calcination temperature dependent Co-Mn

ferrites were prepared by sol-gel combustion method. Iron ferrite (Fe3O4) thin films

were deposited on Si(100) substrates using a high vacuum PLD apparatus. The effect

of annealing temperature and time on Fe3O4

The x-ray diffraction patterns revealed single phase cubic spinel structure of

all the ferrite series, prepared either from ceramic or sol-gel method. However, in case

of Fe

films was investigated. These (soft)

ferrite materials with different compositions and concentrations were studied from

both Physics and material science point of view.

3O4

The effect of Al

thin films (~143 nm) deposited by PLD at room temperature and in-situ

annealed, at 300 to 450 °C temperature. The present work focuses on the effect of

annealing on the thin films and demonstrates the conditions under which increasing

magnetization can be achieved. 3+ contents on the structural, electrical and magnetic

properties of CuZn ferrites was studied by considering the compositions

Cu0.5Zn0.5Fe2-xAlxO4 (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5). The samples were prepared by a

simple and economical, solid state reaction method. It was observed that lattice

constant decreased gradually from 8.385 Å to 8.211 Å with the increase of Al3+

contents, which was attributed to smaller ionic radius of Al3+ as compared to Fe3+.

Temperature dependent DC electrical resistivity decreased with increase in

temperature confirming its semiconductor behavior. The decrease in saturation

magnetization could be understood considering the non-magnetic nature of aluminum.

Dielectric constant, tangent of dielectric loss and loss factor, all showed decreasing

Conclusion

102

trend with increasing frequency confirming high frequency applications of the Al3+

substituted Cu-Zn ferrite samples.

Lattice parameter and crystallite size of Co0.5Mn0.5Fe2O4 nanoparticles

prepared with sol-gel combustion method increased with increasing calcination

temperature. The increase in crystallite size with temperature might be attributed to

the cluster formation of individual crystallites and their increased sizes. Decrease of

coercivity at larger size was due to development of domain walls in nanoparticles. No

significant change was observed in saturation magnetization as a function of

calcination temperature. In our case, both as-burnt and calcined powders are shown to

possess crystalline behavior, contrary to previous results.

Low temperature single phase nanocrystalline Mn0.5Cu0.5-xZnxFe2O4 (x = 0,

0.1, 0.2, 0.3, 0.4, 0.5) and Mn0.5Cu0.5-xNixFe2O4 (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5) ferrites

were also successfully prepared by sol-gel combustion technique. In the first series,

increasing trend of lattice parameters was observed with Zn2+ contents, which was

obviously due to its larger ionic radius as compared to Cu2+ ions. Coercivity

decreased with zinc concentration due to increase in crystallite size. As zinc ions had

strong preference for A-site, transferring Fe3+ ions from A-site to B-site, therefore an

enhanced saturation magnetization was observed as the Zn2+ contents were increased.

The same behavior was observed in case of Ni2+ substituted MnCu ferrites where, the

coercivity was also decreased and saturation magnetization increased in the series.

The trend was attributed to the substitution of ferromagnetic Ni2+ contents in place of

diamagnetic Cu2+ ions.

Ferrite thin films have their own independent and unique importance,

particularly in magnetic storage devices. In this context, iron ferrite thin films having

composition Fe3O4 were deposited on Si(100) substrates by PLD at various

temperatures ranging from room temperature (RT) to 450 °C. The XRD patterns of

the films showed the inverse-spinel structure with different orientations. All samples

showed ferromagnetic behavior but surprisingly we obtained high magnetization of

854 emu/cc at 450 °C, which was higher than the bulk value (471 emu/cc) of Fe3O4

In another work Fe

.

This enhanced magnetization was attributed to iron rich regions within the films.

When the deposition and annealing temperatures was increased, the crystallite size

was also observed to increase but the coercivity was decreased.

3O4 films were deposited by PLD on Si(100) substrates at

450°C and in-situ annealed for 30, 60 and 90 minutes. Here, we obtained [111]

Conclusion

103

oriented films with single phase cubic structure independent of substrate orientation.

By increasing the annealing time, the crystallite size was increased but there was a

decrease in saturation magnetization which might be due to some anti-phase

boundaries between the films and the substrates.

Proposals for Future Work

We have put lots of effort to ensure that all the prepared samples have right

stoichiometry, are single phase and possess high crystalline quality, by using XRD,

SEM and VSM. In the case of Mn-Cu-Zn and Mn-Cu-Ni ferrites, there is insufficient

literature reporting synthesis single phase by sol-gel combustion. We feel that our

present study has been successful in answering some of the questions posed about

these ferrites. However at the same time, there are some points that have evolved and

require further clarification. The following investigations are proposed for future

work.

1. Due to narrow range of stoichiometry of Fe3O4

2. A low temperature magnetic property is necessary to confirm the Verwey

transition temperature of Fe

, XPS or Raman

Spectroscopy is needed to confirm the exact phase of this material.

3O4

3. Depositing Fe

.

3O4

4. Transition electron microscopy is required to understand the structural

changes occurring due to Ni substitution and correlating them with their

magnetic properties.

at room temperature is challenging. The thin film

deposition on various types of substrates using different deposition

techniques are required to gain a comprehensive understanding of their

magnetic properties with respect to the respective deposition techniques.

Appendix

104

Appendix Published Papers 1. Effect of Temperature on Structural and Magnetic Properties of Laser Ablated

Iron Oxide Deposited on Si(100)

Shahid M. Ramay, Saadat A. Siddiqi, M. Sabieh Anwar and S. C. Shin

CHIN. PHYS. LETT. Vol. 26, No. 11(2009) 117504

2 Structural, magnetic and electrical properties of Al3+

S.M. Ramay, Saadat A. Siddiqi, S. Atiq, M.S. Awan, S. Riaz

Chin. J. Chem. Phys. Vol. 23, No. 5 (2010) 591

Accepted Paper 1. Influence of temperature on the structural and magnetic properties of

Co


0.5Mn0.5Fe2O4 ferrites

S.M. Ramay, Saadat A. Siddiqi, S. Atiq, M. Saleem, S. Naseem, M. Sabieh

Anwar, Bulletin of Material Science, 2011

CHIN. PHYS. LETT. Vol. 26,No. 11 (2009) 117504

Effect of Temperature on Structural and Magnetic Properties of Laser AblatedIron Oxide Deposited on Si(100)

Shahid M. Ramay1, Saadat A. Siddiqi1, M. Sabieh Anwar2**, S. C. Shin3

1Centre for Solid State Physics, University of Punjab, Lahore-54590, Pakistan2School of Science and Engineering, Lahore University of Management Sciences (LUMS), Opposite Sector U, D.H.A.

Lahore 54972, Pakistan3Department of Physics, KAIST, 373-1 Guseong-dong, Yuseong-gu, Daejeon, 305-701, Republic of Korea

(Received 20 April 2009)

We fabricate Fe3O4 thin films on Si(100) substrates at different temperatures using pulsed laser deposition, andstudy the effect of annealing and deposition temperature on the structural and magnetic properties of Fe3O4

thin films. Subsequently, the films are characterized by x-ray diffraction (XRD), scanning electron microscopy(SEM) and vibrating sample magnetometery (VSM). The XRD results of these films confirm the presence ofthe Fe3O4 phase and show room-temperature ferromagnetism, as observed with VSM. We demonstrate theoptimized deposition and annealing conditions for an enhanced magnetization of 854 emu/cm3 that is very highwhen compared to the bulk sample.

PACS: 75. 70.−i, 75. 79.−v, 75. 60.−v, 07. 55. Jg, 07. 70.Ds

It is well known that properties of Fe3O4 thinfilms are strongly dependent on crystal structure andgrowth conditions. As far as the magnetic proper-ties of iron based compounds are concerned, Fe3O4

and 𝛾-Fe2O3 are ferrimagnetic, Fe is ferromagnetic,and 𝛼-Fe2O3, FeO are antiferromagnetic.[1] Fe3O4 hasbeen predicted to possess half-metallic properties withhigh spin polarization (100%) of the charge carriersat the Fermi level. Furthermore, it has a relativelyhigh Curie temperature (860 K).[2] Several half metal-lic materials like half Heusler alloys (NiMnSb),[3,4]

full Heusler alloys (Co2MnSi),[5,6] chromium dioxide(CrO2),[7,8,9] pervoskites (La0.7Sr0.3MnO3),[10,11] andmagnetite (Fe3O4)[12,13] are known. Out of thesematerials, Fe3O4 is especially attractive because ofits promising applications in spintronic devices, mag-netic storage, and as a source of spin-polarized currentinjection.[14−18]

There is a growing amount of literature discussingthe optimum growth and deposition conditions ofFe3O4 due to their immense technological importance.A number of deposition techniques have been used togrow Fe3O4 thin films on different substrates, at differ-ent working pressures and from different phases of Fetargets.[19−22] However, it is still difficult to grow themwith well defined compositions and structures at roomtemperature (RT). Pulsed laser deposition (PLD) isconsidered to be one of the best techniques that allowscontrolled film growth at low temperatures. In PLD,formation of oxides is favored by the presence of smallamounts of O2 inside the deposition chamber. Theoxygen while interacting with the ablation plume pro-motes incorporation into the growing film.[23] A fewreports about the annealing effect on the structuraland magnetic properties of Fe3O4 thin films[24−26] ex-ist, but their import and potential still remain unclear.

The present work focuses on the effect of annealing onthe thin films and demonstrates the conditions underwhich increasing magnetization can be achieved.

In this study, we have fabricated Fe3O4 thin filmson Si(100) substrates at different temperatures (fromRT to 450∘C) by pulsed laser deposition, and inves-tigated the effect of annealing and deposition tem-perature on the structural and magnetic properties.Phase analysis of the films was carried out by x-raydiffraction with Cu 𝐾𝛼 radiation. Grain size, lat-tice strain and lattice constants were measured bythe Williamson–Hall plot,[27] assuming that the peakshapes are Lorentzian. Crystal structure determina-tion was performed using the PowderX software.[28]

Film thickness and magnetic properties were deter-mined from scanning electron microscopy (SEM) andvibrating sample magnetometery (VSM) respectively.

The films were deposited on Si(100) substrates byablating a commercially purchased Fe3O4 target hav-ing one inch diameter and 5 mm thickness. Beforedeposition, the Si(100) substrates were subjected tochemical cleaning using acetone and isopropanol in anultrasound bath. The cleaned substrates were thenpre-heated in the deposition chamber at 500∘C for30 min under a vacuum of 1 × 10−7 torr to removeoxides.

The deposition procedure lasted for 20 min. Fordifferent experiments, the substrate was kept at dif-ferent temperatures from RT to 450∘C. In each case,a background pressure of 10−6 torr was achieved af-ter adjusting the oxygen flow rate to 0.2 sccm. Thepulse repetition rate, energy density of the Nd:YAGlaser (Ekspla NL-303) and the target from substratedistance were set at 10 Hz, 1.3 J/cm2 and 35 mm, re-spectively. The laser wavelength was 266 nm.

Subsequent to deposition, film thickness was deter-

**Email: [email protected] 2009 Chinese Physical Society and IOP Publishing Ltd

117504-1

http://cpl.iphy.ac.cn

http://www.cps-net.org.cn

http://www.iop.org


mined using scanning electron microscopy (SEM) (Hi-tachi S-4800). Furthermore, the RT deposited filmswere annealed at 350, 400 and 450∘C for 1 h in basepressure of 1×10−6 torr. The crystalline structure andphases were determined by x-ray diffraction (XRD)(Rigaku D/Max-Rc MPA) using Cu 𝐾𝛼 radiation. Fi-nally, the RT magnetic properties were measured byvibrating sample magnetometry (VSM) (Ricken Den-shi, Japan).

Fig. 1. X-ray diffraction patterns of Fe3O4 thin films onSi(100) substrates deposited at room temperature and an-nealed at the shown temperatures.

Figures 1 and 2 show the 𝜃 − 2𝜃 scan XRD pat-terns of Fe3O4 thin films from RT to 450∘C on Si(100)substrates. The lattice parameters were calculated us-ing PowderX software.[28] The grain size and latticestrains were calculated from the Williamson–Hall plotand the equation,[29]

𝐵 cos 𝜃 = 𝑘𝜆/𝐿 + 𝑆 sin 𝜃,

where 𝐵 represents the full width half maximum, 𝐿 isthe volume averaged grain size, 𝑘 is the Scherrer con-stant, in our case taken to be 1, 𝑆 is the lattice strain,and 𝜆 is the wavelength of Cu 𝐾𝛼 radiation.

As the annealing temperature was raised from300∘C to 450∘C, the diffractograms provide evidenceof single phase polycrystalline Fe3O4 thin films, exceptfor an 𝛼-Fe peak, the peak intensity being approxi-mately 10% of the intensity of the strongest Fe3O4

peak. The crystallinity of the films increases with the

increasing annealing temperature and the peaks be-came sharper, showing, as expected, that the grainsize increases with the increasing annealing tempera-ture.

The peak positions also shift to higher angles ascompared to bulk Fe3O4. This shifting is related touniform lattice strains in the plane of the films[29] com-pared to the bulk sample. The situation was, in fact,different for different deposition temperatures. Theas-deposited films show smaller lattice strains as com-pared to annealed films due to the larger film thick-ness. When the thickness is larger, the strain effectdue to substrate is minimized, explaining the reduc-tion in strain with increasing deposition temperatures.However, the trends in the unit cell and crystallite sizeare identical in the annealed and as-deposited films.The results are presented in Tables 1 and 2. The in-creasing trend for the crystallite size has also beenreported by Tangel et al.[30] and Parames et al.[31]

The increasing trend of the lattice strain with the an-nealing temperature appears counter-intuitive as oneexpects annealing to result in reduced strains. In fact,the strain depends on two parameters, one is the ra-tio of the lattice constants of the substrate and thefilm and the second is the ratio of the coefficients ofthermal expansion (𝛼) for the substrate and the film.In our case, the second factor is the more dominant.The coefficient of thermal expansivity for Fe3O4 is10.4×10−6 K−1 (at 300∘C)[32] and is four times the co-efficient for Si, 2.6×10−6 K−1. This means that Fe3O4

expands more than the Si substrate as the annealingtemperature is increased, resulting in increased latticestrains.

Fig. 2. XRD diffraction patterns of as deposited thinfilms. The deposition temperatures are also shown.

Table 1. XRD and VSM analysis of annealed Fe3O4 thin films on Si(100) substrates.

Annealing Lattice Lattice Estimated crystallite 𝑀𝑠 𝐻𝑐

temperature (∘C) constant 𝑎 (A) strain 𝑆 size (nm) (emu/cm3) (Oe)300 8.375 0.012 26.5 431 422400 8.369 0.017 77.0 649 390450 8.361 0.027 184.9 854 325

117504-2



Table 2. XRD and VSM analysis of as-deposited Fe3O4 thin films on Si(100) substrates.

Depositing Lattice Lattice Estimated crystallite 𝑀𝑠 𝐻𝑐

temperature (∘C) constant 𝑎 (A) strain 𝑆 size (nm) (emu/cm3) (Oe)350 8.369 0.0035 18 374 449400 8.364 0.0011 22 231 318450 8.348 0.0010 30 477 274

-3000 -2000 -1000 0 1000 2000 3000

-1000

-500

0

500

1000

Ms (

em

u/cm

3)

H (Oe)

A

B

C

Fig. 3. In-plane magnetization curves of annealed thinfilms with A, B and C representing the samples annealedat 300, 400 and 450∘C.

-3000 -2000 -1000 0 1000 2000 3000

-600

-400

-200

0

200

400

600

C

A

B

Ms (

em

u/cm

3)

H (Oe)

Fig. 4. In-plane magnetization curves of as-deposited thinfilms with A, B and C representing the deposition temper-atures of 350, 400 and 450∘C.

Fig. 5. Variation of 𝐻𝑐 with crystallite size of annealedthin films.

Fig. 6. Variation of 𝐻𝑐 with crystallite size of as-deposited thin films.

Figures 3 and 4 show the RT. magnetization hys-teresis behavior of annealed and as-deposited samples.The figures clearly show hysteretic behaviour suggest-ing their RT ferromagnetic property. The variations ofcoercivity 𝐻𝑐(Oe) with grain size are shown in Figs. 5and 6. The saturation magnetization 𝑀𝑠 and coer-civity 𝐻𝑐 as a function of annealing and depositingtemperature (for both the annealed and as-depositedsamples) are also summarized in Tables 1 and 2. Inboth the cases the coercivity decreases with an in-creasing annealing temperatue, deposition tempera-ture and crystallite size. At an annealing tempera-ture of 450∘C, we obtain an unusually high satura-tion magnetization 𝑀𝑠 = 854 emu/cm3 as comparedto the bulk magnetization of Fe3O4 single crystals(471 emu/cm3).[33] Kennedy and Stampe have also re-ported a high saturation magnetization for the mag-netite thin films grown on Si(100) substrate.[34] Theseauthors used Fe as the target material and suggestedthat the increasing saturation magnetization could bedue to the increasing Fe content in the thin films,which, due to its amorphous structure, does not ap-pear in their XRD results. In our XRD patterns, sharppeaks for crystalline Fe become visible in the regionof 2𝜃 ≈ 45∘. The enhanced magnetization is verylikely due to the presence of iron-rich centers. Thehysteretic loops, apparently, indicate a single phasemagnetic material but it is possible that the coerciv-ity of the Fe is much too small (as compared to the𝑀 −𝐻 loop step size) to distinctly make an appear-ance in the 𝑀 − 𝐻 loop. In such a case, we cannotexpect the hysteresis loop to confirm the synthesis of asingle phase. Now there is the problem of the origin ofFe in the deposited films. It has been reported[35] thatthe stoichiometry of the iron oxide phase, i.e. the Fe:Oratio, depends on the growth conditions such as the

117504-3



oxygen partial pressure, flow rate, laser fluence andsubstrate temperature. It is very likely that duringdeposition, phases of Fe1−𝑥O (wustite) and hematite(Fe2O3) are also formed, which are reduced to Fe asannealing takes place under vacuum. For example, ithas been demonstrated that FeO is stable only at atemperature greater than 570∘C and is decomposedinto 𝛼-Fe and inverse spinel Fe3O4 below 570∘C.[36]

500 nm

500 nm

117 nm

143 nm

Si(100)

(b)

(a)

Fe3O4

Fe3O4

Si(100)

Fig. 7. SEM images of (a) Fe3O4 thin film annealed at450∘C and (b) as-deposited film at 450∘C.

In summary, we have fabricated Fe3O4 thin filmsby PLD at different temperatures (from room to450∘C) on Si(100) substrates and studied the effect ofannealing and depositing temperatures on the struc-tural and magnetic properties of the Fe3O4 thin films.XRD patterns of both series show the cubic inverse-spinel structure with different orientations. Annealingincreases the crystallinity of the samples. As far as themagnetic properties are concerned, we obtain the fer-romagnetic behavior of all the thin films and obtaina surprisingly high magnetization of 854 emu/cm3 at450∘C (annealed temperature) which is higher thanthe bulk value (471 emu/cm3) of Fe3O4. This may bedue to iron rich regions within the films as alreadyreported. By increasing the annealing and depositingtemperatures, the lattice parameters and coercivitydecrease, while the volume average crystallite size in-creases.

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117504-4


CHINESE JOURNAL OF CHEMICAL PHYSICS VOLUME 23, NUMBER 5 OCTOBER 27, 2010

ARTICLE

Structural, Magnetic, and Electrical Properties of Al3+ SubstitutedCuZn-ferrites

S. M. Ramaya, Saadat A. Siddiqia, S. Atiqb∗, M. S. Awanc, S. Riaza

a. Center of Excellence in Solid State Physics, University of the Punjab, Lahore-54590, Pakistanb. School of Science and Engineering, Lahore University of Management & Sciences (LUMS), Lahore54972, Pakistanc. COMSATS Institute of Information Technology, Islamabad, Pakistan

(Dated: Received on March 10, 2010; Accepted on July 16, 2010)

Nanocrystalline Cu0.5Zn0.5AlxFe2−xO2 (x=0.0, 0.1, 0.2, 0.3, 0.4, and 0.5) ferrite materialswere synthesized using standard solid state reaction technique. The effects of Al3+ contentson the structural, electrical, and magnetic properties were investigated. Single phase cubicspinel structure was revealed by X-ray diffraction analysis. The crystallite size was evaluatedconsidering the most intense diffraction peak (311) using Scherrer formula. Lattice constantdecreased, whereas porosity increased with the increase in Al3+ concentration. The valueof saturation magnetization decreased with increasing aluminum contents. Temperaturedependent value of direct current electrical resistivity has been determined. It is observedthat the substitution of Al3+ has significant impact on the dielectric constant, tangent ofdielectric loss angle and dielectric loss factor. The variation in dielectric properties wasattributed to space charge polarization.

Key words: Oxide material, Ferrite, Solid state reaction, Electrical resistivity, Dielectricconstant

I. INTRODUCTION

Ferrite materials have attracted a considerable atten-tion of the researchers for decades due to their interest-ing soft magnetic properties and high frequency appli-cations [1]. A proper choice of cations along with Fe2+,Fe3+, and their distribution between tetraherdral (A-site) and octahedral (B-site) sites of the spinel lattice,imparts useful and interesting electrical and magneticproperties to the spinel ferrites. Further tailoring ofthese properties using appropriate preparation method,chemical composition, sintering time, and doping ad-ditives always help to improve the technological appli-cability of the ferrite materials [2]. It is essential tocontrol the electrical resistivity of the spinel ferrites inorder to corporate these materials for a wide range ofapplications. This can be achieved in two ways of con-trolling the sintering temperature and choosing properelemental substitution. Excellent dielectric propertiesof ferrites further extend their application range frommicrowave to radio frequencies. The useful frequencyrange is fixed by the onset of resonance phenomenonfor which either the permeability starts to decrease ata critical frequency or the losses rise rapidly [3]. Re-cently, Cu-Zn based ferrites have been synthesized, ex-

∗Author to whom correspondence should be addressed. E-mail:[email protected]

hibiting high Curie temperature with a little compro-mise on initial permeability [3, 4]. The presence of Cuions in ferrites activates the sintering process leadingto increase in density and decrease in losses. While, itis well known that Zn content exerts important influ-ence on the microstructure and hence on the magneticproperties of ferrites. The substitution of Al3+ in fer-rites could lower the dielectric constants that warranttheir applications for high frequency applications, forinstance as micro wave absorbers.

In this work, we have investigated systematically theeffect of Al3+ substitution on the structural, magneticand electrical properties of Cu0.5Zn0.5Fe2O4. The elec-trical behavior of the samples have been discussed incontext of temperature dependent resistivity, and fre-quency dependent dielectric constant (ε′), tangent ofdielectric loss angle (tanδ), and dielectric loss factor(ε′′).

II. EXPERIMENTS

Samples of Cu0.5Zn0.5AlxFe2−xO4 (x=0, 0.1, 0.2, 0.3,0.4, 0.5) ferrites were prepared by the standard solidstate reaction technique using analytical grade reagents.Low cast CuO (99%), ZnO (99%), and Fe2O3 (97%)in their respective stoichiometric ratios were mixed toprepare the ferrite samples. Grinding of every samplewith specific composition was carried out in agate mor-tar and pestle for 4 h. The samples were calcined in

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592 Chin. J. Chem. Phys., Vol. 23, No. 5 S. M. Ramay et al.

TABLE I Lattice constant a, lattice volume V , sintered density ρs, X-ray density ρx, porosity P , saturation magnetizationMs, and activation energy ∆E of Cu0.5Zn0.5AlxFe2−xO2 ferrite system with different Al3+ composition x.

x a/A V /A3 ρs/(g/cm3) ρx/(g/cm3) P Ms/(emu/mL) ∆E/eV

0.0 8.385 589.53 4.59 5.41 0.151 75 0.450

0.1 8.383 589.11 4.38 5.28 0.170 72 0.452

0.2 8.340 580.09 4.27 5.23 0.183 68 0.393

0.3 8.317 575.31 4.15 5.16 0.195 56 0.437

0.4 8.266 564.79 4.10 5.10 0.196 50 0.462

0.5 8.211 553.59 3.96 5.07 0.219 38 0.440

the muffle furnace at 800 C for 8 h. After the in-situcooling of the samples in the furnace, each sample wasground again for 2 h. The samples in powder form werepelletized (diameter of 15 mm) using Apex hydraulicpress by exerting a uniaxial pressure of 4.5×103 kg for3 min. The samples were annealed at 1100 C for 44 hin order to get the required phase.

The investigation of the crystal structure was carriedout using a Rigaku D-Max II-A, diffractometer systemwith Cu Kα (λ=1.5406 A) radiation. Surface morphol-ogy and microstructural features such as grain size andporosity were examined using Hitachi S-3400, scanningelectron microscopy (SEM). The grain size was mea-sured using the line intercept method.

As ferrites are highly resistive materials, thereforetwo-probe method was employed to determine the elec-trical resistivity of the samples in the temperature rangefrom room temperature (RT) to 480 K. Frequency de-pendent (up to 1 MHz) RT measurements of dielec-tric constant and dielectric loss were obtained usinga QuadTech-1920 LCR Meter. Magnetic characteriza-tions were performed using a Lake Shore-7404 vibratingsample magnetometer (VSM).

III. RESULTS AND DISCUSSION

Figure 1 shows X-ray diffraction (XRD) patterns ofthe samples Cu0.5Zn0.5AlxFe2−xO4 (for x=0, 0.1, 0.2,0.3, 0.4, 0.5). As can be seen in Fig.1, all the samplescan be indexed as having a single phase cubic spinelstructure. No impurity peak was noticed. The breadthof the characteristic ferrite peaks is an indication oflower crystallite size of the samples. The crystallite sizewas estimated from the X-ray peak broadening of (311)diffraction peak using the Scherrer formula [5]. For allthe samples, the crystallite size remained in the rangeof 25–30 nm. The values of the lattice constant a ofthe cubic spinel calculated using the CELL software arelisted in Table I. A decrease in lattice constant was ob-served with increase of Al3+ concentration in samples.The decrease in lattice constant is justifiably expectedand can be attributed to the substitution of smallerAl3+ (0.51 A) for large Fe3+ (0.64 A) in the systemCu0.5Zn0.5AlxFe2−xO4. The bulk density (ρb) was cal-culated from the weight and dimensions of the sintered

220 311222 400 422 511 440

x=0.0

x=0.1

x=0.2

x=0.3

x=0.4

x=0.5

2θ / ( )o30 40 50 60 70

FIG. 1 XRD patterns of Cu0.5Zn0.5AlxFe2−xO4 ferrite sam-ples with different Al3+ composition.

samples using the relation, ρb= m/V [6], where m is themass and V is the volume of the samples. As obviouslyseen from the Table I, the value of the bulk densitydecreased from 4.59 g/cm3 to 3.96 g/cm3 as the Al3+concentration increased from x=0.0 to 0.5 in the series.The decrease in bulk density may be due to the fact thatAl has smaller atomic weight (26.98 a.u.) as comparedto Fe (55.85 a.u.). X-ray density (ρx) of the sampleswas calculated using the relation, ρx=8M/Naa3 givenby Smit and Wijn [7], where M is the molecular weightof the samples, Na is the Avogadro’s number and a isthe lattice constant. The number 8 is included in theformula as there are eight molecules per unit cell inthe cubic spinel ferrite structure. The value of ρx de-creased from 5.41 g/cm3 to 5.07 g/cm3 with the increasein Al3+ contents in the sample series as the decreasein mass overtakes the decrease in volume of the unitcell. It is noted that ρx of each sample is greater thanthe corresponding bulk density which is an evidence ofthe presence of pores in the samples. The porosity wasfound to increase from 0.151 to 0.219 in the series whichis direct evidence that the substitution of Al3+ for Fe3+

leaves relatively more empty spaces in the samples.Figure 2 illustrates the representative micrographs of

the Cu0.5Zn0.5AlxFe2−xO4 system that reveal surfacemorphology of the samples obtained using SEM. Theimages show that the grain size increases with increas-ing Al3+ concentration and lies in the range of about

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Chin. J. Chem. Phys., Vol. 23, No. 5 Al3+ Substituted CuZn-ferrites 593

10 µm 10 µm 10 µm

10 µm 10 µm 10 µm

(a) (b) (c)

(d) (e) (f)

FIG. 2 SEM micrographs of Cu0.5Zn0.5AlxFe2−xO4 with with different Al3+ composition. (a) x=0.0, (b) x=0.1, (c) x=0.2,(d) x=0.3, (e) x=0.4, and (f) x=0.5.

FIG. 3 Saturation magnetization Ms plotted against Al3+

concentration x.

2–6 µm. The increased grain size in the series refersto the more porous samples as is evident from the in-creased value of porosity discussed earlier.

The magnetic hysteresis loops for the series of sam-ples were obtained using vibrating sample magnetome-ter. The results revealed that the value of saturationmagnetization Ms decreased with the increase of Al3+concentration as shown in Fig.3. The trend can be un-derstood by the substitution of a non-magnetic element(Al) for a magnetic element (Fe) at the B-site of thecubic spinel structure has caused the magnetization todecrease gradually [8].

Figure 4 shows the temperature dependent variationin direct current (DC) electrical resistivity measuredby two-probe method. The DC electrical resistivity in-creases as the Al3+ concentration increases for all thesamples, which can be due to the conduction mecha-nism in ferrites which takes place mainly through the

FIG. 4 Direct current electrical resistivity lnρ ofCu0.5Zn0.5AlxFe2−xO4 ferrite samples with different Al3+

composition plotted against temperature.

hopping of electrons between Fe2+ and Fe3+ at B-sitesas explained by Vervey et al. [9]. The hopping proba-bility depends upon the separation of ions involved andthe activation energy. As the distance between two met-als ions at B-sites is smaller than the distance betweentwo metal ions, one at A-site and another at B-site,therefore the electron hopping between A and B siteshas a less probability as compared to hopping betweenB-B sites. Hopping between A and B sites does notlimit for the simple reason that there are only Fe3+

at A site and only Fe2+ preferentially occupy B siteduring processing. Therefore, the deficiency of Fe2+

with increasing Al3+ concentration gives further reasonfor the increase of DC electrical resistivity. The mea-sured values of DC electrical resistivity at 293 K werefound to vary from 2.16×106 Ωcm to 1.17×108 Ωcm as


594 Chin. J. Chem. Phys., Vol. 23, No. 5 S. M. Ramay et al.

FIG. 5 Dielectric constant ε′ of Cu0.5Zn0.5AlxFe2−xO4 fer-rite samples with different Al3+ composition plotted againstfrequency.

the concentration of Al3+ increased from x=0 to 0.5.High values of DC electrical resistivity and relativelyeasy preparation method make ferrites an appropriatechoice for the cores of intermediate and high frequencyelectromagnetic absorbers.

The slopes of the linear plots of DC electricalresistivity, shown in Fig.4, determine the activa-tion energy in the measured temperature range. InCu0.5Zn0.5AlxFe2−xO4 system, the values of activationenergy were found to vary from 0.393 eV to 0.462 eV. Inferrites, the activation energy is often associated withthe variation of mobility of charge carriers rather thantheir concentration. This activation energy plays an es-sential role in overcoming the electrical energy barrierexperienced by the electrons during hopping process,which in turn, contributes towards conductivity.

Figure 5 shows the variation of dielectric constant(ε′) with rise of frequency up to 1 MHz. The valueof ε′ is higher at lower frequencies and is found to de-crease with increase in frequency. At high frequencies,particularly for the composition having x=0.3 to 0.5,the value becomes small, constant and independent offrequency [10]. The variation in dielectric constant isdirectly related to space charge polarization. The pres-ence of higher conductivity phases (grains) in the insu-lating matrix (grain boundaries) produces localized ac-cumulation of charge under the influence of an electricfield, results in space charge polarization [11]. A finitetime is needed for the space charge carriers to line uptheir axes parallel to an alternating electric field. Acontinuous increase in field reversal frequency resultsin a point where space charge carriers cannot remainpreserved with the field and the alternation of their di-rection lags behind the field, resulting in a reductionof dielectric constant of the material [12]. In addition,space charge polarization also results from inhomoge-neous dielectric structure of the material as proposed byMaxwell and Wagner in the form of two-layer model [13,14]. According to this model, space charge polarizationoriginates from large well conducting grains separated

FIG. 6 Tangent of dielectric loss angle tanδ ofCu0.5Zn0.5AlxFe2−xO4 ferrite samples with different Al3+

composition plotted against frequency f .

by thin poorly conducting intermediate grain bound-aries. In ferrites, polarization can also be regarded asa similar process to that of conduction [15]. The hop-ping of electron between Fe3+ and Fe2+, results in thelocal displacement of electrons in the direction of ap-plied field that contributes towards polarization. Whenthe frequency is increased, polarization decreases until aconstant value. Beyond this critical value of frequency,the electron exchange between the two cations cannotfollow the alternating field.

Predominance of species like Fe2+, oxygen vacancies,grain boundary defects, and voids contribute signifi-cantly to increase the dielectric constant at lower fre-quencies [16]. At higher frequencies, any species con-tributing to polarizability lags behind the applied fieldand hence the decreasing trend in dielectric constant iswitnessed.

The tangent of dielectric loss angle (tanδ) decreasedwith the increase of frequency as shown in the Fig.6. Itis essential to note that the value of tanδ depends ondifferent factors such as stoichiometry, Fe2+ content andstructural homogeneity. These factors, in turn, dependon the composition of the samples and their sinteringtemperature [17]. The decrease of tanδ with an increasein frequency could be explained on the basis of Koopsphenomenological model [18].

An essential part of the total core loss in ferrites istermed as dielectric loss factor (ε′′) [19]. Figure 7 showsthe plot of frequency dependent dielectric loss factor.As the number of hopping electrons increase, the ex-tent of local displacement in the direction of electricfield increases, causing an increase in electric polariza-tion, which in turn enhances dielectric loss. The dielec-tric losses in ferrites are exhibited during conductiv-ity measurements, as highly conducting materials showhigh losses [20]. Therefore, the present ferrite serieswith relatively low losses might be useful in technolog-ical applications at higher frequencies.


Chin. J. Chem. Phys., Vol. 23, No. 5 Al3+ Substituted CuZn-ferrites 595

FIG. 7 Dielectric loss factor ε′′of Cu0.5Zn0.5AlxFe2−xO4 fer-rite samples with different Al3+ composition plotted againstfrequency f .

IV. CONCLUSION

Aluminum substituted CuZn-Ferrite materials pre-pared by conventional solid state reaction technique ex-hibited single phase, cubic spinel structure, and nano-sized crystallite size. The crystal lattice constant de-clines gradually from 8.385 A to 8.211 A with the in-creasing Al3+ contents. This trend is attributed to thesmaller ionic radius of Al3+ as compared to Fe3+. Thedecrease in DC electric resistivity of all the samples withincreasing temperature depicts the semiconductor likebehavior of the samples. The reason for decrease insaturation magnetization with increasing Al3+ contentsin the CuZn-ferrite series could be understood by thenon-magnetic nature of aluminum. The dielectric con-stant, tangent of dielectric loss and dielectric loss fac-tor showed decreasing trend with increasing frequencyensuring high frequency applications of the Al3+ sub-stituted CuZn-ferrite samples.

V. ACKNOWLEDGMENT

We are grateful to Dr. S. Naseem, Dr. M. S. An-war and M. Saleem for their help in the experimental

measurements and useful discussions.

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