Spin Tronic s

THIN FILM OXIDES AND HETEROSTRUCTURES FOR

SPINTRONICS

Sourav Chattopadhyay

i

THIN FILM OXIDES AND HETEROSTRUCTURES FOR

SPINTRONICS

Thesis submitted to the

Indian Institute of Technology, Kharagpur

For award of the degree

of

Doctor of Philosophy

by Sourav Chattopadhyay

Under the guidance of

Dr. T. K. Nath

& Dr. P. Banerji

DEPARTMENT OF PHYSICS AND METEOROLOGY

INDIAN INSTITUTE OF TECHNOLOGY KHARAGPUR

MAY 2011 © 2011 Sourav Chattopadhyay. All rights reserved.

ii

APPROVAL OF THE VIVA-VOCE BOARD

Certified that the thesis entitled THIN FILM OXIDES AND HETEROSTRUCTURES FOR

SPINTRONICS submitted by SOURAV CHATTOAPDHYAY to the Indian Institute of

Technology, Kharagpur, for the award of the degree Doctor of Philosophy has been accepted by

the external examiners and that the student has successfully defended the thesis in the viva-voce

examination held today.

(Member of the DSC) (Member of the DSC) (Member of the DSC) (Supervisor) (Supervisor) (External Examiner) (Chairman)

iii

CERTIFICATE

This is to certify that the thesis entitled “Thin film oxides and heterostructures for

spintronics”, submitted by Mr. Sourav Chattopadhyay to Indian Institute of

Technology, Kharagpur, is a record of bona fide research work under my supervision and

is worthy of consideration for the award of the degree of Doctor of Philosophy of the

Institute.

__________________________ ______________________ Superviser Superviser

Date: Date:

iv

Acknowledgements

I wish to thank Dr. T. K. Nath (supervisor) for introducing the research field and

discussion about my research work. He has not only supervised my work but was also a

source of constant inspiration. I find him approachable, communicative, open to ideas and

suggestions, and very encouraging. He has prepared the base of my understanding regarding

the experimental techniques and findings. I heartily acknowledge his friendly introduction to

every field including both official and academic. Above all, I acknowledge him for the

independence, presented by him, during my research work that enable me to learn the skill of

work self sufficiently to certain extent. I wish to thank Dr. Pallab Banerji (co-supervisor)

also.

It is a pleasure to thank the members of my Doctoral Scrutiny Committee (DSC), Dr. C

Jacob, Dr. A Dhar and Dr. S. Das for their constant encouragement.

I would like to thank present Head of the Department, Prof. B. K. Mathur and former

Head of the department, Prof. R. N P. Choudhary for providing me the research facility.

I wish to express my deep appreciation to Prof. G. A. Gehring, Prof. A. M. Fox, Dr. A.

J. Behan, Dr. J. R. Neal, D. Score and Q. Feng, University of Sheffield, UK, for their kind

help in arranging several measurements and encouragement and suggestions about my

research work.

I would like to thank Indian Nanoelectronic Users Program, IIT Bombay, for

enormous helping for deposition and growth of device structures in clean room environment

and giving some measurement facilities.

I also would also like to thank all of my lab mates (Sourav Kundu, Samir Kumar Giri,

Pampa Rani Mandal, Proloy Taran Das, Jaganandha Panda and Dhiren Kumar

Prodhan) and seniors (Sanjay Kumar Mandal and Puja De) for their continuous help to

carry out the research work.

I would like to acknowledge Central Research Facility, FIST facility, IIT Kharagpur

for different measurements and DST-NSTI (No. SR/S5/NM-04/2005), India for use of PLD

and FE-SEM experimental facility.

I would like to acknowledge Advance Technology Development Center (ATDC) for

allowing me to measure thickness using ellipsometry.

v

I also would like to thank Mr. Mohanlal Ghosh (Technician, MMM lab) for making the

pressure contact set ups and other electrical measurement set ups and Mr. Kisto Mallik

(Technician, Hall lab) for helping to carry out different kind of works.

I would like to thank DST, India for financial support to buy several measurement

equipments through Project No. - IR/S2/PU-04/2006.

I am thankful to Indian Institute of Technology Kharagpur for financial support during

the course of this study.

Sourav Chattopadhyay

vi

DECLARATION I certify that

a. The work contained in the thesis is original and has been done by myself under the general supervision of my supervisor(s).

b. The work has not been submitted to any other Institute for any degree or diploma. c. I have followed the guidelines provided by the Institute in writing the thesis. d. I have conformed to the norms and guidelines given in the Ethical Code of Conduct

of the Institute. e. Whenever I have used materials (data, theoretical analysis, and text) from other

sources, I have given due credit to them by citing them in the text of the thesis and giving their details in the references.

f. Whenever I have quoted written materials from other sources, I have put them under quotation marks and given due credit to the sources by citing them and giving required details in the references.

Signature of the Student

vii

Curriculum Vitae

Name : Sourav Chattopadhyay

Date of birth : 31st day of December 1979

ACADEMIC CREDENTIALS

Degree University/ Institute Subject Year of passing

M. Tech Jadavpur University Nano Science and Nano Technology

2006

M. Sc. University of Calcutta Electronic Science 2003

B. Sc. University of Calcutta Electronics 2001

ACADEMIC AWARDS

Award of Institute Research Scholarship from Indian Institute of Technology Kharagpur, India on 21st Aug, 2006.

viii

LIST OF PUBLICATIONS

A. Research Papers Published in International Journals

1. Electrical and magnetoelectronic properties of La0.7Sr0.3MnO3/SiO2/p-Si heterostructure for spintronics application, S. Chattopadhyay, P. Dey and T. K. Nath, Current Applied Physics doi:10.1016/j.cap.2011.02.00 (accepted)

2. Enhancement of room temperature ferromagnetism of Fe-doped ZnO epitaxial thin films with Al co-doping, S. Chattopadhyay, T.K. Nath, A.J. Behan, J.R. Neal, D. Score, Q. Feng, A.M. Fox, G.A. Gehring, Journal of Magnetism and Magnetic Materials vol. 323, pp. 1033 (2011)

3. Temperature dependent carrier induced ferromagnetism in Zn(Fe)O and Zn(FeAl)O thin films by S. Chattopadhyay, T.K. Nath, A.J. Behan, J.R. Neal, D. Score, Q. Feng, A.M. Fox, G.A. Gehring Applied Surface Science vol. 257, pp. 381 (2010)

4. Room temperature enhanced positive magnetoresistance in Pt and carrier induced Zn(Fe)O and Zn(Fe,Al)O dilute magnetic semiconductor junction) by S. Chattopadhyay, T. K. Nath Journal of Applied Physics vol. 108, pp. 083904 (2010). Selected for Virtual Journal of Nanoscale Science & Technology for the October 25, 2010.

5. Electrical properties of Pulsed Laser Deposited ZnO thin films by Sourav Chattopadhyay and Tapan Kumar Nath Advanced Materials Research Vol. 67 , pp. 121 (2009)

6. Electrical characterization of p-ZnO/p-Si heterojunction by S. Majumdar, S. Chattopadhyay and P. Banerji Applied surface science vol. 255, pp. 6141 (2009)

7. Tunneling current at the interface of silicon and silicon dioxide partly embedded with silicon nanocrystals in metal oxide semiconductor structures by G. Chakraborty, S. Chattopadhyay, C. K. Sarkar and C. Pramanik Journal of Applied Physics vol. 101, pp. 24315 (2007)

B. Research Papers communicated in International Journals

1. On investigation of origin of junction magnetoresistance in La0.7Sr0.3MnO3/SiO2/p-Si heterostructures, S. Chattopadhyay and T. K. Nath, Journal of Physics D:Applied Physics

2. Enhanced temperature dependent junction magnetoresistance in the heterojunctions with La0.7Sr0.3MnO3 and iron doped ZnO carrier induced dilute magnetic semiconductors by S. Chattopadhyay, J. Panda, T. K. Nath, Journal of Applied Physics.

3. Extraordinary Hall effect, electronic-and Magneto-transport behavior of carrier induced dilute magnetic Zn(Fe)O and Zn(Fe,Al)O thin film by S. Chattopadhyay and T. K. Nath, Physical Review B.

ix

4. Low-temperature resistivity minima in colossal magnetoresistive La0.7Sr0.3MnO3 thin film: A quantum interference effect by S. Chattopadhyay and T. K. Nath, Solid state communications.

B. Papers presented in Conferences/Symposia

1. Temperature dependent anomalous Hall Effects in DMS Zn(Fe,Al)O epitaxial thin film by S. Chattopadhyay and T. K. Nath, 55th DAE Solid State Physics Symposium 2010 (2010).

2. Temperature dependent junction magnetoresistance behavior of LSMO/Zn(Fe,Al)O heterojunction for spintronics by J. Panda, S. Chattopadhyay and T. K. Nath, 55th DAE Solid State Physics Symposium 2010 (2010).

3. J.Panda,S.Chatopadhyay,T.K. Nath,Temperature dependent junction magnetoresistance behavior of the Ni nanoparticle in TiN with p-Si heterojunction,ICONQUEST, NPL, 2010.

4. Investigation on La0.7Ca0.3MnO3/SiO2/n-Si and La0.7Sr0.3MnO3/SiO2/p-Si MOS like heterostructures for Spintronics by S. Chattopadhyay, S. K. Giri and T. K. Nath, International Conference on Fundamental & Applications of Nanoscience and Technology (ICFANT) (2010).

5. Magnetoresistive behavior of epitaxial Zinc oxide thin films doped with iron by S. Chattopadhyay, T. K. Nath International Conference on Magnetic Materials (ICMM-2010) (2010)

6. Room temperature magnetic sensors with Zn(FeAl)O by Pt Schottky contact by S. Chattopadhyay, T. K. Nath 54th DAE Solid State Physics Symposium (2009)

7. Electrical properties of Zn/La0.7Sr0.3MnO3/Pt Schottky device for spintronics by S. Chattopadhyay, T. K. Nath Condensed Matter Days (CMDAYS09) (2009)

8. Electrical properties of La0.7Sr0.3MnO3/SiO2/Si MOS structure by S. Chattopadhyay, P. Dey, T. K. Nath 53rd DAE Solid State Physics Symposium (2008)

9. Electrical properties of Pulsed Laser Deposited ZnO thin films by S. Chattopadhyay, T. K. Nath International Conference on Nanomaterials and Devices Processes and applications (2008)

10. I-V characteristics of La0.7Sr0.3MnO3/SiO2/Si MOS structure by S. Chattopadhyay, P. Dey, T. K. Nath National Seminar on Advanced Nanomaterials and its Applications (2008)

x

Abstract

This work contains the study of the properties of two kinds of spintronics materials,

namely, dilute magnetic semiconductor (DMS) and colossal magnetoresistive (CMR) half

metallic ferromagnetic manganites with very high spin polarization. The DMS materials, namely,

wide band gap Zn(Fe)O and Zn(Fe,Al)O epitaxial films have been chosen with different Fe

concentrations (5, 7 and 10%). The structural (XRD, FESEM, TEM, AFM etc), magnetic

(M(H,T), Anomalous Hall Effect), Optical (UV-VIS absorption spectroscopy down to 5 K),

electrical (resistivity, Hall, Magnetoresistance etc.) properties have been investigated explicitly

and the room temperature carrier induced ferromagnetic behavior have been observed in these

DMS systems. The junction properties of Zn(Fe)O and Zn(Fe,Al)O with Pt have been studied

and all the junction shows positive junction magnetoresistance and this behavior is strictly found

to depend on the magnetic moments of the DMS materials. It can be well described using spin

injection theory. Highly spin polarized, half metallic, ferromagnetic CMR manganites,

La0.7Sr0.3MnO3 thin films have been chosen as a potential spintronic electrode materials and its

structural, magnetic, electronic- and magneto-transport properties have been investigated in

details. Temperature dependent electrical and magneto-transport studies have been carried out on

those films and possible transport models have been examined. The La0.7Sr0.3MnO3/Si/SiO2 MOS

like junctions show positive junction magnetoresistance and it is temperature dependent where

the dominating current transport mechanism through the junctions is found to be Frenkel-Poole

type tunneling. The origin of positive MR has been explicitly investigated for these junctions.

The junction properties of La0.7Sr0.3MnO3 with ZnO, Zn(Fe)O and Zn(Fe,Al)O heterojunctions

have also been studied in details and the junctions show high positive to negative junction

magnetoresistance depending on temperature and magnetic field. The appearance of junction

magnetoresistance in all these Schottky and heterojunctions are best explained using standard

spin injection theory.

Keywords: Semiconductor Spintronics, Dilute magnetic semiconductors, Colossal magnetoresistive manganite, Spin injection, Magnetic heterojunction.

xi

List of Symbols

A* Richardson constant A Area of junction B Magnetic flux D Diffusivity Ec Conduction band Eg Band gap Ev Valence band H Magnetic field Ihkl diffraction intensity of the crystal plane (hkl) of the deposited film

Iohkl diffraction intensity of the crystal plane (hkl) of the bulk standard samples

j ↑ Spin up current density

j↓ Spin down current density J Current density J0 Reverse saturation current density k Boltzmann constant m* Effective mass M Magnetization MRint Intrinsic magnetoresistance MRspt Spin polarized tunneling magnetoresistance

n↑ Spin up electron concentration

n↓ Spin down electron concentarion NA Acceptor ion concentration nc Carrier Concentration ND Donor ion concentration ni Intrinsic carrier concentartions P Spin polarization PjF(0) Current spin polarization q Electronic charge

R↑FM Majority spin up electron

R↓FM Minority spin down electron

R0 Normal Hall co-efficient RAP Anti-parallel resistance

xii

rc Contact resistance rF Ferromagnetic resistence rFN Effective equilibrium resistance rN Non-ferromagnetic resistance RP Parallel resistance Rs Anomalous Hall co-efficient RS Series resistance S Spin T* Cross over temperature T Temperature TM Low temperature minima TP Metal-insulator transition temperature V Voltage V0 Turn on voltage vd Drift velocity Σ Total conductance Σ↓ Spin down conductance Σ↑ Spin up conductance ε0 Vacuum permittivity εs Dielectric constant η Ideality factor θD Debye temperature μ Mobility μF(0) Ferromagnetic sides of the junctions of junction μn(0) Electrochemical potentials for non-magnetic side of junction ρ Resistivity σ Conductivity ΦB Barrier height χdia Diamagnetic susceptibility χpara Paramagnetic susceptibility

xiii

List of abbreviations

AFM Atomic force microscope AHE Anomalous Hall Effect BMP Bound magnetic polaron CIP Current-In-Plane CMR Colossal magnetoresistance CPP Current Perpendicular-to-Plane DE Double-exchange DI De-ionized DMS Dilute magnetic semiconductor DOS Densities of states EDAX Energy dispersive X-ray EVRH Efro’s varieable range hopping F/N Ferromagnet/nonmagnet interfaces FC Field cooled FESEM Field emission scanning electron microscopic FET Field effect transistor FM Ferromagnetic F-N Fowler-Nordheim F-P Frenkel-Poole FWHM Full width at half maximum GMR Giant magnetoresistance HRTEM High resolution transmission electron microscope HRXRD High resolution X-ray diffraction I-V Current-Voltage JMR Junction magnetoresistance J-V Current density-Voltage MR Magnetoresistance MRAM Magnetoresistive random-access-memory MTJ Magnetic tunnel junction NEXAFS Near edge x-ray absorption fine structure NM Nonmagnetic OHE Ordinary Hall Effect PLD Pulsed Laser Deviation RKKY Ruderman-Kittel-Kasuya-Yosida RT Room temperature RTFM Room temperature ferromagnetism

xiv

SCLC Space Charge Limited current SQUID Superconducting quantum interference device TC Texture coefficients 2DEG Two dimensional electron gas TM Transition metal TMR Tunneling magnetoresistance UV-Vis Ultra violet-visible spectroscopy VRH Variable range hopping VTI XAS

Variable temperature insert X-ray absorption spectroscopy

ZFC Zero field cooled

xv

Contents

Title page iCertificate of Approval iiCertificate iiiAcknowledgement ivDeclaration viCurriculum Vitae viiList of Publications viiiAbstract xList of Symbols xiList of Abbreviations xiiiContent xvChapter 1: Introduction and Literature overview 1.1. Introduction 1 1.2. Literature Overview 2 1.2.1. Spintronic materials and devices 2 1.2.1.1. Giant magnetoresistance 2

1.2.1.2. Tunneling magnetoresistance 4 1.2.1.3. Colossal magnetoresistance 7 1.2.1.4. Dilute magnetic semiconductor 11 1.2.1.4.1. Origin of ferromagnetism in DMS 13 1.2.1.5. Organic spintronics 16 1.2.2. Spin transport mechanism 17 1.2.2.1. Spin drift and diffusion 17 1.2.2.2. Spin injection and spin tunneling 18 1.2.2.2.1. Spin injection and spin extraction 19

1.2.2.2.2. Silsbee-Johnson spin-charge coupling 20 1.2.2.2.3. Spin injection into semiconductors 21 1.2.3. Active magneto-electronic devices 23 1.2.3.1. Spin field effect transistor 23 1.2.3.2. Spin diodes 24 1.2.3.3. Spin bipolar transistor 25 1.3. Scope of the thesis 26 References 27Chapter 2: Experimental equipments and techniques 2.1. Introduction 33 2.2. Brief description of used equipments 33

xvi

2.2.1. Thin film deposition unit: Pulsed Laser Deposition (PLD) 33 2.2.2. Characterization equipments 34 2.2.2.1. Structural and surface morphology 34

2.2.2.1.1. High resolution x-ray diffraction technique (HRXRD)

34

2.2.2.1.2. High resolution transmission electron microscopy (HRTEM)

35

2.2.2.1.3. High resolution field emission scanning electron microscopy (FE-SEM)

35

2.2.2.1.4. Energy dispersive x-ray analysis (EDAX) 35

2.2.2.1.5. X-ray absorption spectroscopy (XAS) 36

2.2.2.1.6. Atomic force microscope (AFM) 37 2.2.2.2. Optical characterizations 37 2.2.2.3. Magnetic characterizations 38 2.2.2.4. Electrical characterization 39

2.2.2.4.1. Cryogen free high magnetic field (Superconducting magnet) VTI system

39

2.2.2.4.2. Electrical Measurement Instruments 40 2.2.2.4.3. Temperature readouts and controller Instruments 40 2.3. Brief description of experimental technique 40 2.3.1. Four probe resistivity measurements 41 2.3.2. Hall Effect measurements 42 References 43Chapter 3: Properties of room temperature ferromagnetic Zn(Fe)O and Zn(Fe,Al)O epitaxial thin film 3.1. Introduction 44 3.2 Experimental procedures 45 3.2.1. Preparation of targets 45 3.2.2. Cleaning of substrates 46 3.2.3. Preparation of thin films 46 3.2.4. Characterization of thin films 46 3.3. Results and discussions 47

3.3.1 Chemical properties study 3.3.2. Structural properties

4747

3.3.3. Surface morphology 49 3.3.4. Optical properties 50 3.3.5. Magnetic properties 51 3.3.5.1. Room temperature magnetic properties 51 3.3.5.2. Low temperature magnetic properties 53 3.3.5.3. Carrier dependent ferromagnetism properties 59 3.3.6. Electrical properties 63 3.3.6.1. Electrical transport properties 64

xvii

3.3.6.2. Hall Effect study 67 3.3.6.2.1. Ordinary Hall Effect 68 3.3.6.2.2. Anomalous Hall Effect 71 3.3.6.3. Magnetoresistance properties 73 3.4. Summary 76 References 77Chapter 4: Junction magnetoresistance of Pt/Zn(Fe)O and Pt/Zn(Fe,Al)O metal-dilute magnetic semiconductorjunction 4.1. Introduction 81 4.2. Experimental procedure 81 4.3. Results and discussion 82 4.3.1. Structural properties 82 4.3.2. Magnetic properties 83 4.3.3 Current-voltage characteristics without applied magnetic field 84 4.3.4. Current-voltage characteristics with applied magnetic field 85 4.3.5. Junction magneto-resistance properties 88 4.4. Summary 90 References 90

Chapter 5: Structural, magnetic and electrical behavior of La0.7Sr0.3MnO3 thin films on p-Si 5.1. Introduction 93 5.2. Experimental procedure 94 5.2.1. Preparation of Targets 94 5.2.2. Cleaning of substrates 94 5.2.3. Deposition of La0.7Sr0.3MnO3 film 96 5.3. Results and discussion 96 5.3.1. Structural study 96 5.3.2. Surface morphology 99 5.3.3. Magnetic properties 100 5.3.4. Electrical transport properties 101 5.3.4.1. ρ-T properties without applied magnetic field 101 5.3.4.2. ρ-T properties with applied magnetic field 102 5.3.4.2.1. ρ-T properties lower TM 105 5.3.4.2.2. ρ-T properties above TM 107 5.4. Summary 108 References 109Chapter 6: Junction magnetoresistance study in La0.7Sr0.3MnO3/SiO2/p-Si heterostructures 6.1. Introduction 111

6.2. Experimental procedure 112 6.3 Results and discussion 114 6.3.1 Structural properties 114

xviii

6.3.2. Electrical properties of LSMO/SiO2/p-Si hereostructure without applied magnetic field

114

6.3.2.1. Current-Voltage study using diode characteristics 114 6.3.2.2. Tunneling Characteristics 118

6.3.3. Electrical properties of LSMO/SiO2/p-Si hereostructure with applied magnetic field

120

6.3.3.1. Current-Voltage properties under magnetic field study using diode characteristics

121

6.3.3.2. Tunneling Characteristics under1 T applied magnetic field 124 6.3.4. Junction magnetoresistance properties study 125 6.4. Summary 129 References 130Chapter 7: Electronic-and magneto transport of La0.7Sr0.3MnO3/ZnO, La0.7Sr0.3MnO3/Zn(Fe)O and La0.7Sr0.3MnO3/Zn(Fe,Al)O heterostructures

7.1. Introduction 132 7.2. Experimental Procedure 133 7.2.1. Preparation of target 133 7.2.2. Cleaning of substrate 133 7.2.3. Preparation of heterojunction 133 7.2.4. Characterization of heterostructure 134 7.3. Results and Discussion 134 7.3.1. Structural and surface study 134 7.3.2. Electrical properties study 136 7.3.3. Junction Magnetoresistance properties 141 7.4. Summary 145 References 145Chapter 8: Conclusions 8.1. Conclusions of thesis 147 8.2. Scope of future work 148 8.3. Contribution of thesis 148

Chapter 1

Introduction and Literature overview

Introduction and Literature overview Chapter 1

1

1.1. Introduction

It has been argued with considerable justification that the last half of the 20th

century could be called the microelectronics era. The Moore’s law even starts to run out

of its momentum one day, as the size of individual bits approaches the dimensions of

atoms. This has been called the end of silicon road map. For this reason and also to

enhance multifunctionality of devices, investigators have been eager to exploit another

property of the electron characteristics known as spin. Spin is a purely quantum

phenomenon. Electrons should spin clockwise and anticlockwise directions. Spin

therefore acts as binary logic ‘one’ and ‘zero’. The movement of spin, like flow of

charge, can also carry information among devices. The spin relaxation and spin transport

phenomena are fundamentally important– not only as basic physics questions but also of

their demonstrated value in electronic technology.

In recent years, ‘‘spintronics’’ has been initiated and is progressing outstandingly.

It is an idea to use the spin of electrons in electronic devices for high-speed, high-density,

non-volatile memories and quantum computation in the future. Spintronics is one which

refers normally to phenomena of electrons playing the decisive role. In wider sense

spintronics is a promising research field of electronics. The physical mechanisms of

electronic spin in semiconductors may ultimately lead to multifunctional device based on

photonics, electronics, and magnetic devices [1]. Using the coherent spin phenomena in

semiconductors [2], this may be fundamental for the viewpoint of quantum computation.

The electrical spin injection into semiconductors using both ferromagnetic and

paramagnetic semiconductors, and more recently with Zener tunneling processes are

intended for potential spin based electronics [3-5].

Though the metal spintronics, such as giant magnetoresistance (GMR) systems

have already been used in the computer hard disk read heads memories the

semiconductor spintronics is yet to demonstrate its full potential in computer industries.

Semiconductor spintronics depends on the concepts of spin transport, spin injection, spin

dependent tunneling, as well as spin relaxation and spin dynamics. Spin injection from a

ferromagnetic material into a semiconductor attracts massive attentions to the researchers

in this field. The injection and detection of a spin-polarized current in semiconductors

could combine magnetic storage of information with electronic readout in a single

Chapter 1 Introduction and Literature overview

2

semiconductor device, yielding many obvious advantages. Based on the crystal

symmetries of the materials and the structural properties of semiconductor based

heterostructures, the spin-orbit coupling takes on different functional forms and can give

an effective spin-orbit Hamiltonians in the systems. Most magnetic semiconductor

devices are still theoretical concepts and thus waiting for experimental demonstrations. A

review of selected and few devices is presented.

1.2. Literature Overview

Spintronics can be defined as the art and science of utilizing the spin of the

electron (as well as its charge) to achieve a few ideas [shown in Fig. 1.1]. In a broad

sense spintronics is a study of spin phenomena in solids, in particular metals and

semiconductors and semiconductor heterostructures. Such studies characterize electrical,

optical, and magnetic properties of solids due to the presence of equilibrium and

nonequilibrium spin populations, as well as spin dynamics. These fundamental aspects of

spintronics give us important insights about the nature of spin interactions or spin

exchange couplings in solids. We also learn about the microscopic processes leading to

spin relaxation. The goal of this applied spintronics is to find the effective ways of

controlling electronic properties by spin or magnetic field, as well as of controlling spin

or magnetic properties by electric currents or gate voltages.

1.2.1. Spintronic materials and devices

1.2.1.1. Giant magnetoresistance

The giant magnetoresistance, a beginning of spin electronics, is actually

multilayers of magnetic and non-magnetic metals with individual thicknesses comparable

EElleeccttrroonn ssppiinn

EElleeccttrroonn CChhaarrggee

PPhhoottoonn PPoollaarriizzaattiioonn

SSppiinnttrroonniiccss

Fig. 1.1. The spin based electronics containing both electron spin and electron charge domain.


3

to the mean free paths. The giant magnetoresistance (GMR) effect was discovered at the

end of 80s [6,7]. Investigation of magnetoresistance in thin magnetic multilayers in the

so-called Current-In-Plane (CIP) geometry have revealed a very large change of the

resistance in the antiferromagnetically coupled Fe/Cr multilayers. The effect was much

larger than the observed magnetoresistance in any metallic multilayer before. The same

effect was observed in the so-called Current Perpendicular-to-Plane (CPP) geometry as

shown in Fig 1.2(a) [8]. The fundamental physical phenomenon lying behind such large

change of resistance is the so-called spin valve effect. Fig. 1.2 (b) shows the spin valve

effect in CPP geometry.

The simplest device is metallic multilayer consisting of two ferromagnetic layers

separated by a non-magnetic conductive layer. This layer has ability to change the

metallic interaction between ferromagnetic layers and allows changing their relative

magnetization by an external magnetic field. Such properties can be realized having the

ferromagnetic layers with different coercivity. As the GMR structure consists of non-

magnetic separator in between ferromagnetic layers, it results in the antiferromagnetic

coupling between ferromagnetic layers themselves. If the bias voltage is applied the

electron transport occurs from one ferromagnet to another as shown in Fig. 1.3. In this

case, in the ferromagnetic metal all current is carried by majority spin-up electrons

( ↓↑ < FMFM RR ) and thus is spin-polarized. If the FM/NM interface does not contain large

number of spin scattering, the spin polarized electrons are injected into non-magnetic

layer.

Field

FM

NM

F

Fig. 1.2. Giant magnetoresistance structure in (a) CIP and (b) CPP geometry

FM

FM

NMField

(a) (b)


4

If the layer is thin the spin flips and the spin polarized electrons arrived at second

ferromagnetic interface with preferred spin orientation that backed to the first

ferromagnetic layer. It causes an antiferromagnetic configuration and acts causes high

resistance at the junction. In case of parallel alignment, the current in second

ferromagnetic metal is also carried by spin-up electron and it causes a small junction

resistance. However, the CPP geometry is the easiest for practical realization, since the

resistance of device in CPP geometry is too low to allow direct measurements. A large

number of technological solutions like, superconducting leads [9], sub micron pillars or

rods [10,11] and V-groove [12] have been implemented in order to over come this

limitation.

The typical material combinations in GMR devices are ferromagnetic Fe, Co,

NiFe separated by Cr, Cu, Ag, Au, Re, Ru with typical thickness of ~ 1 to 5 nm. The

magnetic sensitivity can be increased combining a large number of such magnetic

multilayers. These GMR junctions in the relatively week external magnetic fields show

extremely large change of the resistance 220% at low temperatures [13] and 100% at

room temperature [14].

1.2.1.2. Tunneling magnetoresistance

A magnetic tunnel junction (MTJ), which consists of a thin insulating layer sandwiched

between two ferromagnetic electrode layers, shows tunnelling magnetoresistance (TMR)

properties due to spin-dependent electron tunneling through the barrier. Tunneling

magnetoresistance was first reported by Julliere in 1975 [15]. Making with Co–Ge–Fe

sandwich layer Julliere showed the change in electrical resistance with applying a field

and switching the relative alignment of the magnetic moments of Co and Fe from parallel

↑FMR

↓FMR

↑FMR

↓FMR

RN

M FM FM

NM

EF

E

FM

E

FM

E

N

N(E)

(a) (b) Fig. 1.3. Electron spin transport in GMR junction formed by ferromagnetic metal; (a) layered circuit diagram, (b) band diagram for spin injection process in GMR.


5

to anti-parallel directions. He reported a 14% increase in resistance at a temperature of

4.2 K. Julliere’s work may have been inspired in part by the work of Tedrow and

Meservey [16,17] who had earlier measured the spin-dependence of tunneling currents

through an amorphous aluminum oxide tunnel barrier separating various ferromagnetic

electrodes from superconducting aluminum. Tunneling magnetoresistance received much

more attention in later periods. In 1995 Miyazaki et al. [18] and Moodera et al. [19]

reported TMR in excess of 10% at room temperature which was sufficient for making

TMR applicable.

The resistance of a magnetic tunnel junction (MTJ), which consists of a thin

insulating layer (a tunnel barrier) sandwiched between two ferromagnetic (FM) metal

layers (electrodes), depends on the relative magnetic alignment (parallel or antiparallel)

of the electrodes as shown in Fig. 1.4. The resistance R of the junction is lower when the

magnetizations are parallel [Fig. 1.4(a)], and it is higher when the magnetizations are

antiparallel [Fig. 1.4(b)] i.e. APP RR < . This change in resistance with the relative

orientation of the two magnetic layers, called the TMR effect, is one of the most

important phenomena in spintronics. The size of this effect is measured by the fractional

e e e e

FM FM Barrier

e

e

FM FM Barrier

EF EF

(a) (b)

(c) (d)

Fig. 1.4. Typical TMR structure, (a) parallel and (b) anti-parallel alignment of magnetic spins. (c) and (d) are the corresponding conduction band density of state structures for TMR junction.


6

change in resistance, P

P

RR−APR , which is called the magnetoresistance ratio (or MR

ratio). In 2001 first-principle calculations predicted that epitaxial MTJs with a crystalline

magnesium oxide (MgO) tunnel barrier would have MR ratios of over 1000%, and in

2004 MR ratios of about 200% were obtained at RT in MTJs with a crystalline MgO (0 0

1) barrier. The huge TMR effect in MgO-based MTJs is nowcalled the giant TMR effect

and is of great importance not only for device applications but also for clarifying the

physics of spin-dependent tunnelling.

MR ratios of above 200% have recently been observed at room temperature in

fully epitaxial MTJs with MgO (0 0 1) tunnel barrier and Heusler-alloy electrodes [20].

This large TMR effect, however, is thought to originate from the coherent tunnelling in a

crystalline MgO (0 0 1) barrier rather than from the half-metallic nature of the electrodes.

When a crystalline MgO (0 0 1) barrier is used with simple ferromagnetic electrodes such

as bcc Fe, Co and CoFeB yield MTJs with MR ratios from 180% to 500% at RT [21-24].

Sakuraba et al. [25] observed a MR ratio of 570% at low temperature in MTJs with an

amorphous aluminium oxide barrier and Heusler-alloy electrodes. They also observed a

feature characteristic of a spin-dependent tunnelling in those MTJs. This giant TMR

effect at low temperature is therefore thought to be due to the half-metallic nature of

Heusler-alloy electrodes.

The best explanation of TMR effect is proposed by Julliere. This famous paper

proposed a simple phenomenological model, in which the TMR effect is due to spin-

dependent electron tunneling. According to this model, the MR ratio of an MTJ can be

expressed in terms of the spin polarizations P of the ferromagnetic electrodes,

21

21

12

PPPP

MR−

= (1.1)

)()()()(

FF

FF

EDEDEDED

P↓↑

↓↑

+

−=

αα

ααα ; α = 1 and 2. (1.2)

Here αP is the spin polarization of a ferromagnetic electrode, and )( FED ↑α and

)( FED ↓α are, respectively, the densities of states (DOS) of the electrode at the Fermi

energy for the majority-spin and minority-spin bands.


7

Magnetoresistive random-access-memory (MRAM) cells with very large ratios of

parallel to anti-parallel conductance can enable a new type of computer architecture. This

kind of MRAM structure can be achieved by TMR structures. Such kind of devices

would be similar to a field programmable gate array that could be reprogrammed on a

nanosecond timescale. High density MRAM cells (Fig. 1.5 (b)), should have MR ratios

higher than 150% at room temperature, and the read head in the next generation

ultrahigh-density hard disk drive should have both a high MR ratio and an ultra low

tunnelling resistance in TMR structures [26].

1.2.1.3. Colossal magnetoresistance

Half-metallic properties were first discovered by Groot et al. [27] based on band

structure calculations in NiMnSb and PtMnSb crystals. Later the perovskite manganites

doped with alkali metals attracts much attention of researchers due to their half-metallic

behavior with unusual high spin polarized (~ 100%) electronic band structure.

During last decades, numbers of different compounds derived from LaMnO3

inspire researchers due to their Colossal Magnetoresistive (CMR) response to applied

magnetic fields [28-30]. This CMR effect and the correlated degrees of freedom of

PP nn++nn++

MMTTJJ

WWLL

WWrriittee WWLL

RReessiissttaannccee ooff MMTTJJ ((RR))

LLoowweerr lleeaadd

UUppppeerr lleeaadd

CCaapp llaayyeerr

AAFF llaayyeerr SSeeeedd llaayyeerr

FFMM eelleeccttrrooddee ((ffrreeee llaayyeerr))

FFMM eelleeccttrrooddee ((ppiinnnneedd llaayyeerr))

TTuunnnneell BBaarrrriieerr

SSyyFF ssttrruuccttuurree

RRPP

(a)

(b)

(c)

RRAAPP

00

MMRR rraattiioo == ((RRAAPP --RRPP)) //

MMaaggnneettiicc ffiieelldd

(d)

Fig. 1.5(a) Schematic circuit diagram and (b) typical cross-sectional structure of a MRAM cell, (c) typical cross-sectional structure of a MTJ for practical applications, (d) A typical magnetoresistance curve of a MTJ and the definition of MR ratio.

TTuunnnneell bbaarrrriieerr

MMJJTT

WWoorrdd lliinnee ((WWLL))

BBiitt lliinnee ((BBLL))

MMOOSS--FFEETT

FFMM eelleeccttrrooddee

FFMM eelleeccttrrooddee


8

magnetic structure, crystallographic structure and electrical resistivity in CMR materials,

in addition to being of fundamental scientific interest, appears to provide some scope for

engineering more sensitive magnetoresistive response. The ‘colossal’ magnetoresistance

(CMR) rare earth manganites display a fascinating diversity of behaviors including

several forms of magnetic, orbital and charge ordering [31-33]. The materials also exhibit

dramatic variations of physical properties with frequency, temperature, chemical

composition and applied strain, as well as the magnetoresistive properties, which give

them their colloquial name. The particular MR phenomena to be described here are the

gigantic decrease of resistance by application of a magnetic field [29,34-35]. This CMR

effects are observed in manganites sparked a great amount of effort aimed at

understanding the electronic and magnetic properties of these materials. At low

temperatures, optimally hole doped manganites exhibit ferromagnetic metallic or nearly

metallic behavior, while at high temperatures they exhibit a paramagnetic insulating

behavior. In addition to the CMR effect, the manganites have been found to exhibit a

very wide range of exotic and interesting phenomena, including many types of magnetic

ordering, metal-insulator transitions, charge and orbital ordering and pressure induced

phase transitions. It should also be remembered that the manganites belong to the class of

materials where electron correlations are deemed important.

Fig. 1.6. Crystal field splitting of five fold degenerate atomic 3-d levels

Jahn Teller distortion

EJT

eg

3d orbitals

Cubic crystal field splitting

t2g

eg

So = 3/2 Core spin

S = ½ Conduction Electron spin

(xy, yz, zx)

(x2-y2, 3z2-r2)

1 eV


9

CMR materials are compounds of manganese (Mn), oxygen (O) and other

elements. The electrically and magnetically important ion is Mn; the Mn is connected by

oxygen, and the other elements play a role in determining the exact crystal structure and

the charge density of the Mn. The important electronic states are the Mn d-levels. The

manganese (Mn) ion in the CMR manganites is surrounded by the oxygen octahedron. In

free space the d-levels are five-fold degenerate, but in a solid, ‘crystal field’ effects

coming from hybridization and the electrostatic interaction with neighboring ions will

partially or wholly lift the degeneracy. In the ideal perovskite structure the crystal field

has cubic symmetry and splits the d-multiplet into a doublet transforming as the eg

representation of the cubic group Oh and a triplet transforming as the t2g representation as

shown in Fig. 1.6. The lower-lying orbitals, t2g states, are dxy, dyz and dzx, while the

higher-lying ones, eg states, are dx2

-y2 and d3z

2-r

2. The crystal field splitting between the t2g

and eg states is about 1 eV. In the Mn3+ based compounds, the Mn site shows the

electronic configuration of 132 gg et (total spin number S = 2). All the 3d electrons are

subject to electron repulsion interaction or the electron correlation effect. Even the eg

state electrons, hybridized strongly with oxygen 2p states, are strongly affected by such a

correlation effect, and tend to localize in the “carrier undoped” or the parent Mn3+ based

compound, forming the so called Mott insulator. However, the eg electrons can be

itinerant and hence play a role of conduction electrons, when electron vacancies or holes

are created in the eg orbital states of the crystal. The latter hole-doping procedure

corresponds to creation of mobile Mn4+ species on the Mn sites. In contrast, the t2g

electrons, less hybridized with 2p states and stabilized by the crystal field splitting, are

viewed as always localized by the strong correlation effect and as forming the local spin

(S = 3/2) even in the metallic state. The important consequence of the apparent separation

into the spin and charge sectors in the 3d orbital states are the effective strong coupling

between the eg conduction electron spin (S = ½) and t2g localized electron spin (S = 3/2).


10

This on-site ferromagnetic coupling is nothing but the Hund’s rule. The exchange energy

JH (Hund’s rule coupling energy) is as large as 2-3 eV for the manganites and exceeds the

intersite hopping interaction 0ijt of the eg electron between the neighbouring sites, i and j.

In the case of the strong coupling limit )/( ∞→ijH tJ , the effective interaction tij can be

expressed in terms of Anderson-Hasegawa relation,

⎟⎟⎠

⎞⎜⎜⎝

⎛=

2cos0 ij

ijij ttθ

(1.3)

That is the absolute magnitude of the effective hoping depends on the relative angle θij

between the neighbouring (classic) spins. The ferromagnetic interaction via the exchange

of the conduction electron whose spin shows the on-site (Hund’s rule) coupling with the

local spin is called “double-exchange interaction” after the naming by Zener. This

terminology comes from the fact that Zener considered the “double” exchange process of

the electron between the two Mn sites via the oxygen 2p state as shown in Fig. 1.7 By

creating hole doping, the eg electron can hop depending on the relative configuration of

the local spins. The ferromagnetic metallic state is stabilized by maximizing the kinetic

energy of the conduction electrons (θij = 0). When temperature is raised up to near or

above TC, the configuration of the spin is dynamically disordered and accordingly the

Mn3+ Mn4+ tij

θij

t2gLaMnO3 AFM Insulator

egt2g

La1-x SrxMnO3 (T ~Tc)

La1-x SrxMnO3 (T ~Tc)

egt2gH

t2g

La1-x SrxMnO3 (T <<Tc) FM

eg

eg

Fig. 1.7. Schematic diagram of double exchange mechanism


11

effective hopping interaction is also subject to disorder and reduced on average. This

would lead to enhancement of the resistivity near and above TC. Therefore, the large MR

can be expected around TC, since the local spins are relatively easily aligned by an

external field and hence the randomness of the eg hopping interaction is reduced. This is

the simplest explanation of the MR observed for the manganites around TC in terms of

the double-exchange (DE) model [36]. The physics of the colossal magnetoresiatance

(CMR) is obviously more complex. There are other important factors than in the above

simplest DE scenario, e.g. electron-lattice interaction, antiferromagnetic superexchange

interaction between the t2g local spins, inter-site exchange interaction between the eg

orbitals (orbital ordering tendency), intra-site and inter-site Coulomb repulsion

interactions among the eg electrons etc. Among the above interactions other than the DE

interaction, the important electron-lattice interaction stems from the Jahn-Teller type

coupling of the conduction eg electrons with oxygen displacement [37]. The Jahn-Teller

type lattice distortion that lifts the orbital degeneracy and lowers the electronic energy is

frequently observed for the orbital degenerate d-electron configuration. In the crystal,

such a Jahn-Teller distortion is collective and a coherent distortion of metal (e.g. Mn) –

oxygen network is realized, as typically seen in LaMnO3.

Typically doped perovskite oxides with alkali metals are half-metallic in nature

and show good magnetic as well as electronic properties. The Sr doped LaMnO3

manganites or La1-xSrxMnO3 shows Tc above room temperature which drive this CMR

manganite towards technological applications.

1.2.1.4. Dilute magnetic semiconductor

There is an emerging field of semiconductor spin transfer electronics (spintronics)

which aims to utilize the charge carrier spin in dilute magnetic semiconductor.

Ferromagnetic semiconductors are well established materials since long [38]. Some

known ferromagnetic semiconductors are EuS, EuO, CdCr2S4 etc. The main problem

with this materials are there Tc does not cross the temperature over 100 K. The crystal

structures of such materials are quite different and the growth is very difficult. A typical

dilute ferromagnetic semiconductor would consist of a nonmagnetic semiconductor

doped with small amount of transition metals [39-42]. This would hence be known as a


12

dilute magnetic semiconductor (DMS). For the material to be true DMS, its magnetic

dopant spins should retain remanent alignment when influenced by spin polarized free

carriers.

Early studies of DMS materials start with Mn-doped II-VI alloys like (CdTe, ZnS,

HgTe etc) in the 80s decade [43]. The ternary structures of these compounds make them

amendable to tuning the lattice and band parameters by varing alloy composition. The II-

VI compounds are formed by sp3 bonding, incorporating the valance s-electron from

group-II and p-electron from group-VI element. The elemental Mn has half filled 3d shell

and two valance (4s2) electrons. Mn replaces the group-II element by Mn2+. Since the 3d-

shell of Mn is half filled, it requires considerable energy to add an electron. The magnetic

properties of theses alloys are directed by exchange interactions between local atomic

moment and sp-band electrons. In early 90s, the technological advancement in DMS

materials occurred with discovery of ferromagnetism in Mn doped InAs [44,45]. After

that the DMS properties have been found in other III-V semiconductors also.

Unfortunately, the highest Tc reported for GaAs was 110 K [46]. Later GaP [47], GaN

[48-50], AlN [51,52] showed room temperature ferromagnetism.

The main problem with the DMS investigated at this point is clearly the Tc. A

theoretical paper by Dietl et al. [53] calculated that manganese doped semiconductors had

ability to be ferromagnetic at room temperature. The theory is based on the concept that

how carriers in association with localized spins can make it long range ferromagnetic

interactions in a DMS. The localized spins are Mn2+ spins, of the d5 configuration, and

the carriers are holes that originate from shallow acceptors. The interaction is

parameterized by the p-d exchange term which is in exchange energy n0, where n0 is the

total cation site number density and is p-d exchange integral of the system. When Mn

spins are aligned there is an energy difference between the carriers and Mn spins caused

by magnetic moment. This energy difference will lower the decreasing temperature until

they are equal at Tc. An equation for Tc of a system is then obtained by equating the two

energies. The formula shows that high value of p-d exchange integral is required to

achieve the high value of Tc. The data in Fig. 1.8 show the calculated Tc for various

semiconductors with 5% Mn doping and hole concentration 3.5 × 1020 cm-3. This


13

interesting work encourages huge efforts to achieve room temperature ferromagnetism

and better understanding of the systems.

The main disadvantages of DMS in III-V semiconductors are the solubility of

transition metal ion in it. In wide band gap semiconductors still there is a controversy

whether the ferrogmanetism arrises from the secondary impurity phase or not. After the

acceptance of ZnO as a II-VI semiconductor with Wurtzite structure and wide band gap,

the transition metal doped ZnO has been well studied as a dilute magnetic semiconductor.

The interest in ZnO was originally prompted by theoretical predictions concerning hole

mediated magnetism though the experimental work has been almost entirely concerned

with n-type materials, which raises important and interesting scientific issues concerning

the carrier-mediated magnetism. Except Mn, there are several reports on room

temperature ferromagnetism in ZnO doped with other transition metals like Fe, Ni, Cu

etc. also [54-57]. ZnO doped with rare earth element like V, Gd etc. also shows

ferromagnetism at room temperature [58].

1.2.1.4.1. Origin of ferromagnetism in DMS

Understanding the physical mechanism behind magnetic ordering in DMS

materials is an essential ingredient to their further development. However, there is an

incomplete understanding of the origin of ferromagnetism in transition metal doped

semiconductors. There are some theories which is used to describe the ferromagnetism in

the DMS systems.

Fig 1.8. Calculated Curie temperature values for various p-type semiconductors with the hole concentration of 3.5 × 1020 and 5% Mn.

Curie temperature (K)


14

Dietl’s mean field theory: The model assumes that the ferromagnetic exchange

interactions occur between localized spin doped into the semiconductor matrix and are

mediated by charge carriers. This spins are assumed randomly oriented through out the

semiconductors. As shown in Fig. 1.9 the localized spins are aligned with the interaction

with free carrier and causes ferromagnetism in the system [59].

First principle design: Sato and Katayama have employed first principles design to

investigate the appearance of ferromagnetism in both semiconductor and oxide

spintronics [60,61]. The magnetic stability was calculated using density functional theory

within the frame work of local density approximation. Their results were consistent with

Dietl’s theory in case of Mn doping. Their work also pointed about the contribution of d

state at the Fermi level.

Ferromagnetism in a localized carrier regime: In this proposed model ferromagnetism

in the localized spins can be originated from localized carrier. Ferromagnetism in the

localized carrier regime can be explained through the formation of bound magnetic

polarons (BMP). A BMP is a quasi-particle comprised of the localized carrier and the

magnetic atoms encompassed within its radius as shown in Fig 1.10. The localized

carriers are bound to its associated defects. The exchange between the bound carrier and

the magnetic moments tend to align to parallel moment of another inside the BMP. With

lower temperature the radius of BMP grows and starts to overlap to each other. The

overlapping BMPs become correlated and create a long range ferromagnetic ordering

[62,63].

hh++ MMnn++22 MMnn++22

ssiittee ii ssiittee jj

JJ ss((ii)) ss((jj))

Fig.1.9. Magnetic exchange between two Mn ions mediated by delocalized hole


15

Ferromagnetism in spin-split conduction band: Coey et al. [64] have proposed a model

for appearing of ferromagnetism in ZnO like DMS semiconducting materials based on

the spin-split donor impurity band. In this model, the donor defect (i.e. Oxygen vacancy

etc) overlaps on large concentration to form an impurity band. This impurity band can

interact with local magnetic moment through bound magnetic polarons (BMP) and

creates a long range ferromagnetic interaction.

Free carrier mediated ferromagnetism: In Zener mean field approximation, the

inclination of the ferromagnetic alignment of d electron spins is due to the spin coupling

between the incomplete d shell and conduction electron (or hole). Due to the negligible

roaming of magnetic electron and the quantum oscillations of the electron spin

polarization around the localized spins, this model was ultimately abandoned. Dietl et al.

[53] pointed out on this model that, for semiconductor, the effect of quantum oscillations

averages out to zero since the mean distance between the carriers is greater than that

between spins and hence the Zener mean field model becomes equivalent to Ruderman-

Kittel-Kasuya-Yosida (RKKY) interaction model. Considering this model, high carrier

IIssoollaatteedd iioonn IIssoollaatteedd BBMMPPss

OOvveerrllaappppiinngg BBMMPPss

AAnnttiiffeerrrroommaaggnneettiicc ppaaiirr

Fig.1.10. Illustration of bound magnetic polaron


16

density was shown to drive paramagnetic-ferromagnetic phase transition in DMS

materials [65].

Polaron Percolation model: The polaron percolation model tells that when the

concentration of carriers is much smaller than the magnetic impurity, exchange

interaction between the localized carriers and magnetic impurities lead to their mutual

polarization [66]. Due to this interaction BMP is formed and with decreasing temperature

the radius grows and forms a ferromagnetic ordering in DMS.

1.2.1.5. Organic spintronics

Organic spintronics is a new and promising research field where organic materials

are applied to mediate or control a spin-polarized signal. It is hence a fusion of organic

electronics and spin electronics. Organic materials, on the one hand, open the way to

cheap, low-weight, mechanically flexible, chemically interactive, and bottom-up

fabricated electronics. Phenomena in organic semiconductors seem considerably more

complicated than in their inorganic semiconductors. In particular, the characterization

techniques that have proved so successful for inorganic spin electronics cannot be used

for organic materials. Tris-8-hydroxy-quinoline aluminium (Alq3) sandwiched between

transition metal and La0.7Sr0.3MnO3 half metal, establish a clear correlation between spin-

polarization loss in the organic material and the spin-valve signal [67,68].

The n-alkane-dithiolate and 1,4-n-phenyl-dithiolate molecules shows large

magnetoresistance in both the tunnelling and metallic regime. In the case of nickel

contacts the first molecules show tunnelling behaviour with the spin-polarization of the

current mainly given by surface states at the interface between the nickel and the

molecule as shown in Fig. 1.11 [69].

Nickel Sulphur Carbon Hydrogen

Nickel Sulphur Carbon Hydrogen

Fig. 1.11. Structural and electronic properties of (a) Ni(001)/octane/Ni(001) and (b) Ni(001)/tricene/Ni(001) spin-valve.


17

In contrast, in 1,4-n-phenyl-dithiolate the transport is by means of states extending

across the whole molecule, which determine the spin-polarization of the junction. There

have been several investigations of spin-transport through organic molecules. These

include carbon nanotube spin valves [70], electron coherent spin transfer across

molecular bridges [71], spin injection in π-conjugated molecules [72,73] and organic

tunneling junctions [74]. Although these works demonstrate convincingly that spin-

polarized currents can be injected into organic materials with reasonably high efficiency,

there is a general lack of control over the magnetic response of the devices.

1.2.2. Spin transport mechanism

1.2.2.1. Spin drift and diffusion

The total number of electrons is assumed to be preserved and if the electron

densities are ↑n and ↓n for the spin up and spin down states, the total electron density is,

↓↑ += nnn while the spin density is, ↓↑ −= nns

Considering the spin flip probability, 1<<w over a length l (shown in Fig. 1.12), which is

justified for the conduction electrons, one can easily employ the balance equation using

Taylor expansion,

)(2

2

↓↑↑↑↑ −−

∂

∂−

∂

∂=

∂

∂nnw

xn

vx

nD

tn

d (1.4)

)(2

2

↑↓↓↓↓ −−

∂

∂−

∂

∂=

∂

∂nnw

xn

vx

nD

tn

d (1.5)

Adding the two equations the drift-diffusion equation for the density n, can be written as,

sd

sxsv

xsD

ts

τ−

∂∂

−∂

∂=

∂∂

2

2 where

ττw

s

21= (1.6)

P+ P‐

wP+ wP-

x-l x X+Fig. 1.12. Random walk scheme with indicated spin-flip probabilities


18

sτ is the spin relaxation time. Writing the spin drift-diffusion equation in terms of

mobility and employing the continuity equation one can easily get the spin continuity

equation as,

s

s sxj

ts

τ−=

∂∂

+∂∂ (1.7)

Where, xsDeseEejJ ss ∂∂

+=−= μ is the spin current density and μ is the electron spin

mobility. The right hand side represents the spin relaxation. The spin in a given volume

can decrease either by spin current flowing away from the volume, or by spin relaxation.

The current spin polarization can be expressed as,

jj

jjj

P sj =

−= ↓↑ (1.8)

1.2.2.2. Spin injection and spin tunneling

First spin injection model has been proposed by Aronov in 1976 [75]. The

thermodynamics of spin injection has been developed by Johnson and Silsbee for spin

transport across ferromagnet/nonmagnet (F/N) interfaces [76,77]. The theory of spin

injection has been further developed by several researchers [78-83]. In the following

treatments, the formulations of the spin injection problems by Johnson-Silsbee and

Rashba are discussed. Our goal is to find the current spin polarization, )0(jFP , which

determines the spin accumulation, )0(sNμ , in the normal conductor. We will assume that

the lengths of the ferromagnet and the nonmagnetic regions are greater than the

corresponding spin diffusion lengths. The spin injection scheme is exemplified in Fig.

1.13 assuming that the nonequilibrium spin vanishes at the far ends of the junction.

F C N

>>Ls >>LsNContact

x

Fig. 1.13. Scheme of our spin-injection geometry; The ferromagnetic conductor (F) forms a junction with the nonmagnetic conductor (N). The contact region (C) is assumed to be infinitely narrow, forming the discontinuity at x = 0. It is assumed that the physical widths of the conductors are greater than the corresponding spin diffusion lengths.


19

1.2.2.2.1. Spin injection and spin extraction

As shown in Fig. 1.13 there are three distinct regime in ferromagnetic / nonmagnetic

junction, i.e. ferromagnetic layer with length sFL , non magneric layer with length sNL and

contact. The )0(jFP at the ferromagnetic regime can be expressed as;

F

sFFjF Rj

PP)0(1)0(

μσ += (1.9)

RF is an effective resistance that appears in the spin-polarized transport and is roughly

equal to the actual resistance of the region of length sFL . The spin accumulation )0(sNμ at

the non magnetic regime can be expressed as:

NjNsN RjP )0()0( −=μ (1.10)

The Spin accumulation is proportional to the spin current which pumps the spin into the

system. RN is the effective resistance. The greater is the spin diffusion length, the greater

is the spin accumulation. The advantage of the quasi-chemical potential model over

continuous drift-diffusion equations for charge and spin current, is in describing the spin-

polarized transport across the contact region at x = 0. Employing this equation one can

write the spin polarization at the contact:

C

sjC Rj

PP)0(1 μΔ

+= Σ (1.11)

where,Σ

Σ−Σ= ↓↑

ΣP and . ↓↑ Σ+Σ=Σ

↑Σ and ↓Σ are the conductance of spin up and spin down electrons, respectively and

↓↑ΣΣΣ

=4CR . (1.12)

To solve these three equations of spin polarization electrons one needs to assume the

condition that jCjNjFj PPPP === )0()0( (1.13)

The above equalities are justified if spin-flip scattering can be neglected in the contact.

Using the spin current continuity equations, we can solve our algebraic system and

readily obtain for the spin injection efficiency,

NCF

CFFj RRR

PRPRP

+++

= Σσ . (1.14)


20

The spin injection efficiency is the averaged conductivity spin polarization over the three

regions, weighted by the effective resistances. Using the spin accumulation equation in

non magnetic regime, if j < 0, so that electrons flow from F to N, the spin accumulation is

positive, 0)0( >sNμ ; it is spin injection. If j > 0, the electrons flow from N to F,

and 0)0( <sNμ ; it is called spin extraction. If we look at the density of spin polarization,

Pn = s/n, we get for the density of spin polarization in the nonmagnetic region,

jN

NN

sNn Pn

gjeR

ng

eP −== )0()0( μ (1.15)

Since the injected spin polarization is proportional to the charge current, the electrical

spin injection is an example of spin pumping.

1.2.2.2.2. Silsbee-Johnson spin-charge coupling

In electrical spin injection we drive spin-polarized electrons from a ferromagnet

into a nonmagnetic conductor. Nonequilibrium spin accumulates in the nonmagnetic

conductor. The opposite is also true. If a spin accumulation is generated in a nonmagnetic

conductor that is in proximity of a ferromagnet, a current flows in a closed circuit, or an

electromotive force (emf) appears in an open circuit (shown in Fig.1.14). This inverse

effect is called the Silsbee-Johnson spin-charge coupling. This coupling was first

proposed by Silsbee (1980) and experimentally demonstrated by Johnson and Silsbee

(1985) in the first electrical spin injection experiment.

Considering an F/N junction with a special boundary condition: a nonequilibrium spin is

maintained at the far right boundary of the nonmagnetic conductor, one can write

0)( ≠∞sNμ . Accordingly, at the far left boundary of the ferromagnetic region, the spin is

assumed to be in equilibrium, i.e. 0)( =−∞sFμ . The emf is )()( −∞−∞ sFsN μμ . One of our

V

Spin detectionSpin injection

Fig. 1.14. The Johnson-Silsbee non-local spin injection and detection scheme. Spin injected through one F/N junction. The spin detection is done by a different F/N junction, by the Silsbee-Johnson spin charge coupling. Spin diffusion from the injector is indicated by the different shades of grey.


21

goals is to find the spatial profile of the spin accumulation inside the junction. The e.m.f.

represents the drop of the quasi-chemical potential, μ, across the junction. If such a drop

is present, the system acts as a battery: by closing the circuit, charge current flows. In

electrical and spin equilibrium, the quasichemical potential drop must vanish.

From the drift-diffusion model, since j = 0 the integrating of the equation in the F region,

from −1 to 0, and putting 0)1( =−sFμ , one can wirte )0()0()( sFFFF P μμμ σ=−−∞ . Similarly,

for the N region 0)0()( =−∞ NN μμ . There is a drop of the quasi chemical potential in the F

region, due to the spin-polarization of the conductivity, while the quasi chemical potential

is constant over the N region. The sμ can be expressed as,

[ ] sNLxsNsNsNsN ex /)()0()()( −∞−+∞= μμμμ (1.16)

The above equation gives, [ ])()0(1)0( ∞−−=∇ sNsNsN

sN Lμμμ . (1.17)

Using the condition of j = 0, and assuming again that the spin is conserved across

the interface at x = 0, i.e. )0()0( sNscsFs jjjj === , one can obtain the following set of

equations for the spin currents at x = 0; )0()0( sFF

gFFC

RPRPR

μφ ⎟⎟⎠

⎞⎜⎜⎝

⎛ +=Δ Σ and the

quasichemical potential can be obtained, )()()0( ∞<∞++

= sNsNNCF

FsF RRR

Rμμμ .

This allows writing the spin current at the contact as,

)(1)0( ∞++

= sNNCF

s RRRj μ (1.18)

The electrostatic potential drop across the contact is due to the spin polarization of

the ferromagnet as well as due to the spin filtering effects of the contact. The emf can be

developed if an equilibrium spin (Pj) is in electrically contact with a nonequilibrium spin.

This effect allows detection of nonequilibrium spin, by putting a ferromagnetic electrode

over the region of spin accumulation. By measuring the emf across this junction, we

obtain information about the spin in the nonmagnetic conductor.

1.2.2.2.3. Spin injection into semiconductors

In contrast to normal metals and superconductors, creating a substantial current

polarization jP by direct electrical spin injection from a metallic ferromagnet into a


22

semiconductor proved to be more difficult [84-86]. The conductivity mismatch problem

has been demonstrated by Schmidt et al. [87]. Even in the absence of the resistive

contacts, effective spin injection into a semiconductor can be achieved if the resistance

mismatch is reduced by using for spin injectors either a magnetic semiconductor or a

highly spin-polarized ferromagnet. For spin injection in non-degenerate semiconductors,

there can be large effects due to built-in fields and deviation from local charge neutrality

and space charge region. Interfaces making up a semiconductor often develop a space-

charge region. Typical examples are the Schottky contact and the depletion layer in p-n

junctions. Microscopic studies of spin-polarized transport and spin-resolved tunneling

through space-charge regions are still limited in scope. The difficulty lies in the need to

consider self consistently simultaneous charge accumulation and electric field generation,

both affecting transport. Non-self-consistent analyses of a Schottky-barrier spin injection

were performed by Albrecht and Smith [88] and Prins et al. [89], while Osipov and

Bratkovsky proposed an efficient spin injection method using a δ-doped Schottky contact

[90].

The system is depicted in Fig.1.15. The p-n junction has a magnetic n region with a net

equilibrium electron spin nnP 0 , where n stands for the n region. Holes are assumed to be

unpolarized. At small biases, in which the injected carrier density through the depletion

region is still smaller than the equilibrium carrier density, there is no spin injection. Only

with bias increasing to the high injection limit, the spin is injected. The following formula

was obtained for spin injection [91],

( ) ( )( ) ( )R

nL

nR

nR

n

Rn

Ln

Rn

Rn

Ln

Ln

PPPP

PPPPPP

002

0

002

00

1

11

−+−

−+⎥⎦⎤

⎢⎣⎡ −

=δ

δ (1.19)

n region P region Space charge

+ _

X=-l X=0 X=l

Fig. 1.15. Schematic diagram of spin injection through space charge region in magnetic p-n junction


23

where L (left) and R (right) label the edges of the space-charge (depletion) region of a p-n

junction. Correspondingly, RnPδ represents the nonequilibrium electron polarization,

evaluated at R, arising from a spin source.

1.2.3. Active magneto-electronic devices

The spin valve and the magnetic tunnel junction involve ferromagnetic and non

magnetic metal films with or without insulating tunnel barrier. They are compatible with

CMOS technology. But the devices are passive and are not capable to power gain. The

passive devices are adequate for memory applications if the output voltages are

sufficiently large. An active device which have power gain, are of bigger utility and

figure the spine of semiconductor electronics. Recently researches are focused on

integrate spintronics directly with semiconductors by incorporating a semiconductor

spintronics materials in a device structures. This can be able to develop a spintronics

device with power gain and therefore they will be capable to maintain a large fan out

which is necessary to form a high density electronic logic applications.

1.2.3.1. Spin field effect transistor

One attempt to semiconductor spintronics involves a spintronics device and an

application of spin injection theory to semiconducting channel of a field effect transistor

(FET) had proposed by Datta and Das. In Datta-Das structure [92], a ferromagnetic

source and a drain were connected by 2D electron gas channel (2DEG) with a fixed

source to drain distance (Lx) as shown as schematic diagram in Fig 1.16. A

magnetization of both source and drain were oriented along the axis of the channel and an

internal electric field (E) was perpendicular to the 2DEG plane. Carriers were injected at

the source with their spin axis (along x axes) proceed under the applied magnetic field.

By applying a gate voltage to the channel, the internal electric field (E), the effective

magnetic field (H*) and spin phase (φ) can be varied. Increase of gate voltage sweeps the

magnitude of H* to values that causes spin precessions of multiples of π and 2π, and

thereby causes a periodic source to drain conductance. In the past few years, much

research has been carried out involving the spin injection FET using ferromagnetic

semiconductors for spin injection [93], but, unfortunately, these materials cannot be able


24

to show the characteristics necessary for the device applications because of their low

Curie temperatures.

Though the spin injection FET has not yet been realized, progress had been made

and significant problems relating to device applications are understood.

1.2.3.2. Spin diodes

To understand the mechanism of spin injection through the complex heterostructures it is

more convincing to start with magnetic p-n junction. Consider the fact that the electrons

are spin polarized, not holes. Fig.1.17 shows the p-n junction of different magnetic

semiconductor junctions. Spin injection can happen from magnetic n-side [Fig. 1.17(a)]

and spin extraction are expected to happen to the magnetic p-side [Fig. 1.17(b)]. An

external field causes spin splitting of the magnetic n-region and spin up subband is more

populated with carriers. At a low bias (below built-in potential) there is no spin injection.

While there are more spin up carrier in n, the barrier for crossing the space charge region

is exponentially larger for spin up electrons. These two exponential factors cancelled and

there is no net spin injection. As the bias voltage increases, the barrier for crossing the

space charge region reduced and the spin injection become larger. The same effect

pursues the analogous reasoning for spin extraction from a magnetic p-region. Figure

1.17(c) and (d) shows another mechanism for larger magnetoresistance. If non-

equilibrium population of is created in the n-side, the opposing external factors are

minimized and the large spin injection can be expected.

InGaAs

InGaAs

2DEG

Fe contact Fe contact

Gate Source Drain

Fig. 1.16. A schematic diagram of Dutta-Das field effect transistor


25

1.2.3.3. Spin bipolar transistor

The more complex and interesting device is the bipolar junction transistor with

non-magnetic n/magnetic p/non-magnetic n emitter-base-collector configuration as

shown in Fig. 1.18 [94]. Forward bias is applied to base to emitter to lowering the base-

emitter barrier for electron. Reverse bias is applied to base to collector which increases

the barrier for electron transport from base to collector. A population of nonequilibrium

spin is maintained at the emitter using circularly polarized light or spin injection from

ferromagnetic electrode. Nonequilibrium spin is maintained through the emitter-base

narrow depletion layer, and it causes nonequilibrium spin at magnetic base region. This

causes a spin split in base which depends on external magnetic field. Carrier

recombination in the base is negligible and the base current is formed by holes flow to the

emitter. On the other hand, collector current depends on the electron spin injection from

emitter to base, and then to collector. Increasing the external field increases the spin

splitting and nonequilibrium spin electron to the base increases. It results a sensitive

current gain in the bipolar junction transistor under applied magnetic field.

PP nnoonnmmaaggnneettiicc

nn nnoonnmmaaggnneettiicc

PP nnoonnmmaaggnneettiicc nn

mmaaggnneettiicc

PP mmaaggnneettiicc nn

ssppiinn--ppoollaarriizzeedd

PP nn NN

X = XP Xn LFig.1.17. Band diagram for magnetic p-n junction; (a) electrons from magnetic n-region, (b) electrons from magnetic p-region, (c) spin injection extraction through spin polarized n to magnetic p-region, (d) The scheme where the spin is injected from magnetic heterostructure N into the non magnetic n-region which forms a p-n junction in with magnetic p-region

(a) (b)

(c) (d)


26

1.3. Scope of the thesis

Though the metal spintronics, such as giant magnetoresistance (GMR), tunneling

magnetoresistance (TMR) etc. systems have already been used in the computer hard disk,

read heads memories, spin valves, sensor applications and other technological

applications in computer industries, the semiconductor spintronics is still questionable

according to both scientific and technological field. The integrated spintronics allied with

semiconductors can be able to lead the semiconductor industries towards a new era.

In this thesis, an attempt has been made to find out the properties of dilute

magnetic semiconductors (ZnO doped with Fe and sometime co-doped with Al) and half-

metallic highly spin polarized ferromagnetic mangnaites (La0.7Sr0.3MnO3) having Curie

temperature above room temperature. Moreover, these materials have been used to

fabricate magnetic heterojunctions and metal-semiconductor junctions for possible

spintronics applications. The temperature dependent spin injection process through the

junctions by applying varying magnetic field has been focused in this thesis. Chapter-1

discusses about the brief overview of spintronics materials and devices. Chapter-2

describes in brief the experimental details and different equipments used for

characterization. In chapter-3, the structural, magnetic, electronic transport,

Fig.1.18. Magnetic bipolar junction transistor with magnetic base (a) Schematic diagram, (b) band diagram

(a)

(b)

emitter collectorbase

forwar reverse

je jbjc

N P N

we wb

wc

Spin up electronSpin down electron hole


27

magnetotransport, Hall Effect etc. properties of iron doped ZnO and in some cases co-

doped with Al have been discussed in details. The high crystalline quality epitaxial Fe

doped ZnO dilute magnetic semiconductor thin films deposited on (0001) c-plane single

crystalline sapphire substrates show room temperature ferromagnetic behavior with

carrier mediated ferromagnetism properties where the majority carrier is electron. The

spin injection through the Pt and Fe doped and Al co-doped ZnO junctions have been

estimated from the magnetic field dependent current voltage behavior in chapter-4. The

appearance of positive junction magnetoresistance and dependence of magnetoresistance

on the magnetic moment of ZnO doped with Fe have been explained using

Ferromagnetic/Non-magnetic spin injection theory. Chapter-5 discusses about the

structural, magnetic and electronic properties of non epitaxial La0.7Sr0.3MnO3 thin films

on (100) p-Si substrate. The films show room temperature ferromagnetism with colossal

magnetoresistive behavior. In chapter-6, a detailed study of electrical transport

mechanism through the La0.7Sr0.3MnO3/SiO2/p-Si heterojunction with different type of

SiO2 layers and appearance of junction magnetoresistance have been carried out. The

dependency of junction magnetoresistance on trap charges, leakage currents and defects

in SiO2 have been estimated in this chapter. Chaper-7 deals with the p-n heterojuction

formed using p-type La0.7Sr0.3MnO3 half-metallic ferromagnet and n-type Fe doped ZnO

dilute magnetic semiconductor. The spin injection theory through magnetic

semiconductor p-n junction has been employed to describe the junction

magnetoresistance of such heterostructures.

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32

[76] M. Johnson and R. H. Silsbee, Thermodynamic analysis of interfacial transport and of the thermomagnetoelectric system, Phys. Rev. B 35, 4959 (1987). [77] M. Johnson and R. H. Silsbee, Coupling of electronic charge and spin at a ferromagnetic-paramagnetic metal interface, Phys. Rev. B 37, 5312 (1988). [78] P. C. van Son, H. van Kempen, and P. Wyder, Boundary Resistance of the Ferromagnetic-Nonferromagnetic Metal Interface, Phys. Rev. Lett. 58, 2271 (1987). [79] T. Valet and A. Fert, Theory of the perpendicular magnetoresistance in magnetic multilayers, Phys. Rev. B 48, 7099 (1993). [80] A. Fert and H. Jaffrès, Conditions for efficient spin injection from a ferromagnetic metal into a semiconductor, Phys. Rev. B 64, 184420 (2001). [81] S. Hershfield and H. L. Zhao, Charge and spin transport through a metallic ferromagnetic-paramagnetic-ferromagnetic junction, Phys. Rev. B 56, 3296 (1997). [82] G. Schmidt, D. Ferrand, L. W. Molenkamp, A. T. Filip and B. J. van Wees, Fundamental obstacle for electrical spin injection from a ferromagnetic metal into a diffusive semiconductor, Phys. Rev. B 62, R4790 (2000). [83] E. I. Rashba, Theory of electrical spin injection: Tunnel contacts as a solution of the conductivity mismatch problem, Phys. Rev. B 62, R16267 (2000). [84] P. R. Hammar, B. R. Bennett, M. J. Yang, and M. Johnson, Observation of Spin Injection at a Ferromagnet-Semiconductor Interface, Phys. Rev. Lett. 83, 203 (1999). [85] A. T. Filip, B. H. Hoving, F. J. Jedema, and B. J. van Wees, B. Dutta and S. Borghs, Experimental search for the electrical spin injection in a semiconductor, Phys. Rev. B 62, 9996 (2000). [86] H. J. Zhu, M. Ramsteiner, H. Kostial, M. Wassermeier, H. P. Schönherr, and K. H. Ploog, Room-Temperature Spin Injection from Fe into GaAs, Phys. Rev. Lett. 87, 016601 (2001). [87] G. Schmidt, D. Ferrand, L. W. Molenkamp, A. T. Filip and B. J. van Wees, Fundamental obstacle for electrical spin injection from a ferromagnetic metal into a diffusive semiconductor, Phys. Rev. B 62, R4790 (2000). [88] J. D. Albrecht and D. L. Smith, Electron spin injection at a Schottky contact, Phys. Rev. B 66, 113303 (2002). [89] M W J Prins, H van Kempen, H van Leuken, R A de Groot, W Van Roy and J De Boeck, Spin-dependent transport in metal/semiconductor tunnel junctions, J. Phys.: Cond. Mat. 7, 9447 (1995). [90] V. V. Osipov, and A. M. Bratkovsky, Efficient nonlinear room-temperature spin tunneling-emission in ferromagnetsemiconductor heterostructures with extended penetration depth,’’ cond-mat/0307030 (2003). [91] J. Fabian, Igor Žutić and S. Das Sarma, Theory of spin-polarized bipolar transport in magnetic p-n junctions, Phys. Rev. B 66, 165301 (2002). [92] S. Datta and B. Das, Electronic analog of the electro‐optic modulator, Appl. Phys. Lett. 56, 665 (1990). [93] Y. Ohno, D. K. Young, B. Beschoten, F. Matsukura, H. Ohno, D. D. Awschalom, Electrical spin injection in a ferromagnetic semiconductor heterostructure, Nature 402, 790 (1999). [94] J. Fabian, I. Žutić, and S. Das Sarma, Magnetic bipolar transistor, Appl. Phys. Lett. 84, 85 (2004).

Chapter 2

Experimental equipments and techniques

Experimental equipments and techniques Chapter 2

33

2.1. Introduction

In this chapter, we have discussed about detailed experimental techniques and some

major equipments which have been used to carry out our work on oxide thin films and

heterostructures. First we have used pulsed laser deposition (PLD) unit for depositing the thin

films and heterojunctions. The structural and surface morphological characterizations have been

carried out using high resolution x-ray diffraction (HRXRD) technique, high resolution

transmission electron microscope (HRTEM), high resolution field emission scanning electron

microscope (FESEM), energy-dispersive x-ray spectroscopy (EDAX) and near edge x-ray

absorption fine structure (NEXAFS); the optical properties have been studied using UV-VIS

spectrophotometer and magnetic characterizations have been done using superconducting

quantum interference device (SQUID). The electronic-transport, magneto-electronic, Hall Effect

and magneto-transport properties have been investigated using cryogen free high magnetic field

low temperature VTI system with closed cycle helium refrigeration compressor unit. Current and

voltage source-meter have been used for the current – voltage (I-V), resistivity [ρ(T,H)] and Hall

resistivity [ρH(T,H)] etc. characterizations.

2.2. Brief description of used equipments

2.2.1. Thin film deposition unit: Pulsed Laser Deposition (PLD)

The photograph of experimental set up of the PLD system for thin film oxide films and

heterostuctures deposition has been shown in Fig. 2.1.

Fig. 2.1. The experimental set-up of PLD chamber

Chapter 2 Experimental equipments and techniques

34

A. Laser System: COMPexPro™ 201, High-Pulse-Energy KrF Excimer Laser

manufactured by Coherent, Inc. 5100 Patrick Henry Drive, Santa Clara, CA 95054 has been

used for pulsed laser. The wavelength, Pulse Energy and Maximum Average Power of the

laser source are 248 nm (KrF), 700 mJ and 7 W, respectively. Maximum Repetition Rate is

10 Hz.

B. Deposition Chamber: Deposition chamber with software control unit has been

made by Excel Instruments, Mumbai - 93. For substrate temperature we have used PID

Temperature Controller, Dynamic Control System and substrate heater for a maximum

temperature of 850 °C.

C. Vacuum Components: We have used a turbo molecular pump, TMH/TMU 261,

and a rotary vane pump DUO 10/ MC made by Pfeiffer Vacuum Technology AG;

Headquarters/Germany to evacuate the thin film deposition unit to obtain very high vacuum

(~ 10-7 Torr) and to control the oxygen pressure while depositing the various films.

2.2.2. Characterization equipments

Mainly, the structural, surface morphology, optical, magnetic and electrical characterizations have been carried out for all our thin film samples and heterostructures. A brief description of all the techniques used has been presented here.

2.2.2.1. Structural and surface morphology

2.2.2.1.1. High resolution x-ray diffraction technique (HRXRD)

Structural characterizations of our thin films and heterostructures have been carried out

using high resolution x-ray diffractometer (Model: Philips, PW-1729) with monochromatic Cu-Kα

radiation at room temperature. The tube voltage and current have been kept at 40 kV and 30 mA

respectively. The wavelength of the Kα line, which is 1.541Å in this present case, is basically the

weighted average of the wavelengths of its components Kα1and Kα2, Kα1 being twice the wave length

of Kα2 [1],

541.1)554.1540.12(31

=+× (2.1)

In order to obtain as closely monochromatic Kα1 radiation as possible, Ni filter was used to absorb

undesirable Kβ component. During measurements, the resolution of the instrument i.e. 2

1

α

α

KK is 0.5


35

and accuracy of 2θ value is ±0.03o.

2.2.2.1.2. High resolution transmission electron microscopy (HRTEM)

High resolution transmission electron microscopy (HRTEM) images of our films and

heterostructures were recorded employing JEOL, lEM-2010 ultra - high resolution (UHR)

microscope using a LaB6 filament. During experiment the instrument was operated with an

accelerating voltage of electron, E=200 kV. Corresponding relativistic wavelength of electrons

depends on this accelerating voltage E and its value can be obtained using the modified De

Broglie wavelength [2],

[ ] 2/1200 2/12 cmeEeEm

h+

=λ Å (2.2)

where, h is Planck's constant, m0 the rest mass, e the charge of the electron and c the velocity

of light. Thus obtained λ corresponding to E=200 kV, is 0.025 Å. For low magnification

bright field image, this instrument can resolve a minimum dimension of 2 nm of the specimen

under study and the minimum diameter of electron beam can be ~ 20 nm.

2.2.2.1.3. High resolution field emission scanning electron microscopy (FE-SEM)

High resolution field emission scanning electron microscopy (FE-SEM) has been

done using Carl Zeiss SMT Ltd. SUPRATM 40 [Emitter: Thermal field emission type,

Standard Detectors : High efficiency In-lens detector, Everhart-Thomley Secondary Electron

Detector]. The chamber pressure was maintained at ~ 10-5 mbar and gun pressure at~ 10-5

mbar. During experiment the instrument was operated with an accelerating voltage of

electron, E=5.28 kV. Corresponding relativistic wavelength of electrons, as obtained using

Eq. 2.2, is 0.168 Å. This instrument can resolve a minimum dimension of ~ 1 nm at our

working accelerating voltage of E=5.28 kV. We have used gold coated pelletized samples as

specimens for FE-SEM study.

2.2.2.1.4. Energy dispersive x-ray analysis (EDAX)

Energy dispersive x-ray analysis (EDAX) unit of Oxford instruments is attached with

high resolution field emission scanning electron microscope (Carl Zeiss SMT Ltd

.SUPRATM40). During x-ray analysis of our specimen, the working distance was maintained at

15 mm, the chamber pressure at ~ 10-5 mbar and gun pressure at ~ 10-9 mbar. Both point EDAX


36

and bulk EDAX were performed on our samples, depending, upon specific requirement. We

have employed INCA EDS hardware and INCA software, which provides a stable microanalysis

platform. This EDAX unit promises a < 1 eV shift in peak position and resolution between count

rates of 1 kcps and 10 kcps in microanalysis of our samples.

2.2.2.1.5. X-ray absorption spectroscopy (XAS)

If a high energy x-ray (0.1-100 eV photon energy) excites an electron from core level of

an atom, the resultant photoelectron will jump into unoccupied higher energy states. The created

core hole filled either via an Auger process or by capture of electron from another state which is

then followed by the fluorescent photon. There are three main regions found on a spectrum

generated by XAS data (Figure 2). The dominant feature is called the "rising edge", and is

sometimes referred to as X-ray Absorption Near-Edge Structure (XANES) or Near-edge X-ray

Absorption Fine Structure (NEXAFS). The pre-edge region is at energies lower than the rising

edge. The fluorescent photon or Auger electron which is inelastically scattered photoelectron is

been measured to obtain NEXAFS spectra as shown in Fig. 2.2.

NEXAFS spectra are usually measured either through the fluorescent, in which emitted

photons are monitored, or total electron yield, in which neutralization current is monitored [3].

The XAS spectra are measured for a solid sample with some standard and a comparative study

gives the present states of the element in the system.

x‐ray

Photoelectron

VB

n’

n

n’

n

Fluorescent photon

VB

Auger electron

VB

n’

n

Fig. 2.2. The processes of NEXAFS spectra: (a) photo absorption of an x-ray into a core level followed by photoelectron emission, followed by either (b) filling of the core hole by an electron in another level, accompanied by fluorescence; or (c) filling of the core hole by an electron in another level followed by emission of an Auger electron.

(a) (b) (c)


37

2.2.2.1.6. Atomic force microscope (AFM)

The atomic force microscope (naming by Scanning Probe Microscope (SPM) [Model:

Multiview- 1000TM by Nanonics Imaging Ltd. Malcha Jerusalem 91487 Israel] has been used to

characterized the surface morphology of the thin films. It consists of 70 μm AFM/ NSOM

Scanner, 200 nm×200 nm STM Scanner, Varian Turbo molecular pump and Ion pump, Nd:

YAG Laser for NSOM, Normal Si AFM and also optical fiber tip, Liquid Cell Accessories,

Avalanche Photo Diode, Leica microscope etc. An atomically sharp tip is scanned over a surface

with feedback mechanisms to maintain the tip at a constant force (contact mode), or at constant

oscillation amplitude (non-contact & tapping mode). A laser is focused to the back of the

reflective cantilever. As the tip scans the surface of the sample, moving up and down with the

topographical feature of the surface, the laser beam is deflected into a multi-sectioned PSD

which measures the difference in light intensities and their incident positions to measure the

height of sample surface at that position.

2.2.2.2. Optical characterizations

For optical characterization, absorbance spectra of the nanocrystalline sample are

recorded using UV - visible spectrophotometer (Micro pack, DH-2000, Deuterium Halogen

Light Sources) combine the continuous spectrum of an RF-excited deuterium UV-Visual source

and a halogen VIS-NIR light source in a single optical path (fiber optical path). The combined

spectrum sources produce stable spectral output from ~ 200 - 1200 mm in a compact package).

We have done the global correction of measuring instrument using the bulk ZnO sample with

Integration Time: 5000 msec Average: 10, Box car: 10 and Flash Delay: 100.The global

correction is done through the reflection mode placing the beam of the light at the perpendicular

or 60° angle with the sample surface. After this we have recorded absorbance spectra of all

samples fixing the globalize conditions.

Optical transmission spectra of the thin films are recorded at room temperature in an

energy variation of 1 - 4 eV. The band gaps of all the DMS thin films are estimated from the

measured spectra. The optical absorption measurements are carried out on a large number of

samples of various thicknesses. A steep rise in the absorbance near the absorption edge hints a

direct type transition. In a crystalline material with polycrystalline structure both direct or


38

indirect optical transitions are possible depending on the band structure of the material.

Assuming parabolic bands, the relation between α and Eg for the direct transition is given by,

ngEhh )( −∝ γγα (2.3)

and for indirect transition by

)/exp(1)(

1)/exp()(

TEEhB

TEEhA

hD

nPg

D

nPg

θν

θν

να−−

−−+

−

+−= (2.4)

where, Ep is the phonon energy assisting the transition, θD the Debye temperature and are

constants: For a direct transition n = 1/2 or 3/2 depending on whether the transition is allowed

or forbidden in quantum mechanical sense. Similarly, n = 2 or 3 for indirect allowed and

forbidden transition, respectively. The usual method of determining band gap is to plot a graph

between nh /1)( να and νh and look for that value of n which gives best linear graph in the band

edge region.

2.2.2.3. Magnetic characterizations

Magnetic measurements [Magnetization (M) as a function of magnetic field (H) and

Magnetization (M) as a function of temperature (T)] have been carried out using Quantum

Design superconducting quantum interferometer device, commonly known as SQUID

magnetometer, in the dc magnetic field range of 0 ± 55 kOe and in the temperature range of 2 -

330 K.

The Quantum Design MPMS SQUID VSM Ever-Cool system is an integrated pulse-tube

cryocooler system. This eliminates the need to use any liquid cryogens for the operation of the

MPMS SQUID VSM. The SQUID VSM utilizes a 7 Tesla, superconducting, helium-cooled

magnet and accomplishes rapid switching between charging and discharging states and stable

fields with a unique superconducting switching element called the Quick Switch, which changes,

between superconducting and normal states in less than one second. This allows rapid collection

of high precision data. Typical M-H loop up to 5 T would take ~ 60 mins and M-T measurement

in the temperature range of 4 – 300 K would take ~70 mins. Temperature Accuracy of the

SQUID is lesser of ±1% or 0.5 K


39

2.2.2.4. Electrical characterization

The electrical characterizations of all our oxides thin films and heterostructures have

been done mainly using the high magnetic field (8 T) cryogen free superconducting magnet

with variable temperature insert (VTI) system which is operated down to 2 K temperature along

with other devices, e.g., Keithley-2182 nanovoltmeter, Keithley-2612 source-meter (with 1µV

resolution) and Keithley-6221 AC and DC current source. Temperature has been controlled using

Lakeshore (Model 331) temperature controller with the temperature stability better than ± 50 mK.

2.2.2.4.1. Cryogen free high magnetic field (Superconducting magnet) VTI system

The photograph of Cryogen free high magnetic field VTI system set up has been shown

in Fig. 2.3 (a) and (b). Figure 2.3 (c) shows the schematic VTI circuit and the probe that has been

used for different electrical measurements has been shown in Fig. 2.3 (d).

The cryogen free high field measurement system combines the latest cryogen free

technology with sophisticated measurement techniques. It is comprised of the following main

components: (1) Cryo-cooler system with compressor, (2) Cryostat and Magnet, (3) Variable

temperature insert (VTI), (3) Electronics rack with measuring devices, (4) Measurement System

Software and (5) Water chiller with compressor for the cooling of the Cryo-cooler compressor.

The cryo-cooler system is the Gifford McMahon (GM) cryocooler which has the

advantage of greater thermodynamic efficiency and reliable operation in any orientation. It uses

(a) (c)

(b)

(d)

Fig. 2.3. (a) and (b) The photograph of Cryogen free high magnetic field VTI system set up. (c) The schematic VTI circuit and (d) the probe that has been used for different electrical measurements.


40

the compressor to drive moving pistons with regenerators. This cryocooler can provide more

than 50 watts of cooling power on the 60 K stage (the first stage) and up to 1.5 watts of cooling

power at 4 K stage (the second stage). The main function of the first stage is to cool the radiation

shield around the low temperature parts of the system. Cooling for the magnet and the VTI is

provided by the second stage as shown in Fig. 2.3 (c).

2.2.2.4.2. Electrical Measurement Instruments

The electrical characterizations was made mainly using

(i) Keithley 2182 nanovoltmeter: Measure voltage from 1nV to 120V (channel 1); 10nV to

12V (channel 2) with 6 and ½ digit display.

(ii) Keithley 2612 source-meter: Maximum output power and source/sink limits to 30.603 W

per channel maximum. ±20.2 V at ±1.515 A, ±202 V at ±101 mA, four quadrant source

or sink operation. Voltage regulation is 0.01% of range. Load: ±(0.01% of range + 100

μV).

(iii)Keithley 6221AC and DC current source: Current ranges from 2 nA to 100 mA with 0.1

to 0.4% accuracy.

2.2.2.4.3. Temperature readouts and controller Instruments

Lakeshore (Model 331) temperature controller has been used to control and measure the

temperature. It is Proportional-Integral-Derivative type temperature controller (PID) with 0 to

1000 with 0.1 setting (proportional), 1 to 1000 (1000/s) (integral) with 0.1 setting and 1 to 200%

(derivative) with 1% resolution. The sensor is used here is diode (Silicon, GaAlAs Most

thermocouple types RTDs: 100 Ω) Platinum, Platinum, Germanium, Carbon-Glass, Cernox™,

and Rox™ sensor with 50 Ω heater load power.

All the electrical measuring systems are automated using the LABVIEW software

(version 8.5) through GPIB interfacing cables with a PC.

2.3. Brief description of experimental technique

Mainly the electrical characterizations such as resistivity, Hall Effect, magnetoresistance

and current-voltage measurements have been done on oxides thin films and heterostructures in

the entire thesis work along with some optical and magnetic measurements. Brief descriptions of


41

such measuring techniques have been presented here. In order to make these electrical

measurements, low resistance contacts should be made between the films and the connecting

wires in the sample holder. This is done by pressing tiny In (indium) piece onto the sample

corners and then the fine Cu wires are bonded into this for Fe and Fe,Al doped ZnO thin films

grown epitaxially on sapphire substrates. Wire bonding for La0.7Sr0.3MnO3 samples have been

done connecting Ag and Cu wires. Wires bonding for p-Si have been done using Al which have

been deposited by thermal evaporation method. Then the Cu wires have been connected on Al

with Ag paste.

2.3.1. Four probe resistivity measurements

The purpose of the 4-point probe is to measure the resistivity and magnetoresistance of the

films of any semiconductor materials. It can measure either bulk or thin film specimen, each of

which consists of a different expression. The 4-point probe setup for bulk or thick films and thin

films has been shown in Fig. 2.4 (a) and (b).

For bulk or thick film samples where the sample thickness (t) >>spacing between two probes (s) we

can assume a spherical projection of current coming to sample and differential resistance is,

⎟⎠⎞

⎜⎝⎛=Δ

dAdxR ρ (2.5)

If we carry out the integrations between the spaces where voltage is measured (s), we can write,

I

V

Fig. 2.3. (a) Four probe method for measuring resistivity; (b) van der Pauw geometry for measuring resistivity of thin films.

I

V

(a) (b)


42

πρ

πρ

221

2

2

1

2 sxdxR

x

x

== ∫ (2.6)

where probe spacing is uniformly s. Due to the superposition of current at the outer two tips, R =

V/2I. Thus, we arrive at the expression for bulk resistivity,

⎟⎠⎞

⎜⎝⎛=

IVsπρ 2 (2.7)

Thin film measurements are generally done using the Van der Pauw geometry as shown in Fig. 2.4.

(b). For a very thin layer (thickness t <<s) we get current rings instead of spheres. The derivation is

as follows,

2ln22

2

1txt

dxRx

x πρ

π== ∫ (2.8)

Consequently, for R = V/2I, the resistivity for a thin sheet is,

⎟⎠⎞

⎜⎝⎛=

IVt

2lnπρ (2.9)

2.3.2. Hall Effect measurements

When the current flows along the x-axis, with a magnetic field (B) applied in the y axis,

electron starts to drift from current line due to Lorentz force, Bev×− , v is drift velocity. The basic

Hall geometry has been shown in Fig. 2.4 (a). The electrons are therefore deflected from traveling

line (x-axis) and create a potential drop across the surface along z-axis, which is called Hall voltage

(VH).

y

z

xj

B

V = VH

V = 0

I

VFig. 2.4. (a) Hall Effect geometry in a semiconductor bar; (b) van der Pauw geometry for measuring the Hall Effect in thin films

(a) (b)


43

For the steady state, we can write following equations for the x and y directions, with e the

electric charge and n the carrier concentration,

xxe eEvm =τ/ (2.10)

)(0 BvEe xy −−= (2.11)

BvE xy = with )/( nejv xx −= and thus,

)/( neBjE xy −= (2.12)

Combining those equations with yH EE = and xjj = yields

neRH

1−= (2.13)

RH is Hall coefficient. The measurement is very complex in case of measuring the thin films and

it can be done by van der Pauw geometry as shown in Fig. 2.4 (b). A series of resistances in the

crossed geometry altering voltage and current directions can be made as a function of magnetic

field. The measured resistances are then multiplied by the film thickness and plotted with

magnetic field. The gradient of the plot gives the Hall coefficient for all the thin film samples.

References

[1] B.D. Cullity Elements of x-ray diffraction, Addison-Wesley Publishing Company Inc., 2nd Edition, California (1978). [2] G. Thomas and M. J. Joringe, Transmission Electron Microscopy of materials, Wiley-Interscience Publication, John Wiley & sons, New York (1979). [3] http://en.wikipedia.org/wiki/XANES.

Chapter 3

Properties of room temperature ferromagnetic Zn(Fe)O

and Zn(Fe,Al)O epitaxial thin film

This chapter is based on

International journals 1. Enhancement of room temperature ferromagnetism of Fe‐doped ZnO epitaxial thin films with Al codoping, T.K. Nath, A.J.

Behan, J.R. Neal, D. Score, Q. Feng, A.M. Fox, G.A. Gehring, Journal of Magnetism and Magnetic Materials vol. 323, pp. 1033 (2011)

2. Temperature dependent carrier induced ferromagnetism in Zn(Fe)O and Zn(FeAl)O thin films by S. Chattopadhyay, T.K. Nath, A.J. Behan, J.R. Neal, D. Score, Q. Feng, A.M. Fox, G.A. Gehring, Applied Surface Science vol. 257, pp. 381 (2010)

3. Extraordinary Hall effect, electronic‐and Magneto‐transport behavior of carrier induced dilute magnetic Zn(Fe)O and Zn(Fe,Al)O thin film by S. Chattopadhyay and T. K. Nath, Physical Review B. (communicated)

Conferences/Symposia 1. Temperature dependent anomalous Hall Effects in DMS Zn(Fe,Al)O epitaxial thin film by S. Chattopadhyay and T. K. Nath,

55th DAE Solid State Physics Symposium 2010 (2010) 2. Magnetoresistive behavior of epitaxial Zinc oxide thin films doped with iron by S. Chattopadhyay, T. K. Nath International

Conference on Magnetic Materials (ICMM‐ 2010) (2010)

Properties of room temperature ferromagnetic Zn(Fe)O and Zn(Fe,Al)O epitaxial thin film

Chapter 3

44

3.1. Introduction

There has been a great interest in transition-metal-doped diluted magnetic

semiconductors (DMS), which exploit both the spin and the charge of carriers, because

the combination of two degrees of freedom promises new functionality of memory,

detectors, light-emitting sources and possible use in next generation spintronic devices,

e.g., spin-valve transistors, spin light-emitting diodes, non-volatile storage and logic

devices. A theoretical study by Dietl et al. [1] has predicted that some of the manganese

doped semiconductors have the ability to be ferromagnetic at or above room temperature

(e.g. GaN and ZnO). The low magnetic ordering temperature in most of the DMS

materials limits the potential spintronic device applications at room temperature. There is

an ongoing quest for ferromagnetic DMS with Curie temperatures, Tc, far in excess of

300 K for the second generation of spin electronics, as well as a search for transparent

ferromagnets which could add an optoelectronic dimension. It is also possible to control

the ferromagnetic interactions between the localized spins by the carriers [2-5], as well as

the demonstration of efficient spin injections into normal semiconductors [6, 7]. In spite

of several progresses on transition metal doped ZnO diluted magnetic semiconductor as a

spintronic material, much controversy remains concerning the mechanism that causes the

ferromagnetism. The carrier-induced ferromagnetism has been observed in different III-V

[8-10] and II-VI [11-14] semiconductors. The interest in ZnO was originally prompted by

theoretical predictions concerning hole mediated magnetism. However, the experimental

work has been almost entirely concerned with n-type materials, which raises important

and interesting scientific issues concerning the carrier-mediated magnetism. DMS are

mixed spin-fermion systems, involving randomly distributed localized magnetic and

mobile carriers in the semiconductor band. With carrier concentration much smaller than

the magnetic impurity concentration, the DMS systems provide a complimentary limit to

Kondo systems. The coupling between localized impurity spin (S) and mobile valence

band holes can be represented by the exchange interaction – J S.σ, where σ is the fermion

spin operator. A theoretical approach has been made by S. Singh et al. [15] to explain the

carrier induced ferromagnetism in DMS system by using diluted Hubbard model. The

origin of FM in these materials is still an issue of debate. In the currently accepted picture

Chapter 3 Properties of room temperature ferromagnetic Zn(Fe)O and Zn(Fe,Al)O epitaxial thin film

45

for DMS the presence of carriers is essential to mediate the interaction between the

magnetic ions. An investigation with additional dopants to enhance or induce

ferromagnetism in these DMS materials will be an interesting approach to establish the

theoretical picture of carrier induced FM.

Some efforts have been made with various doping element to enhance both the

electrical conductivity and ferromagnetism. In some literature the coexistence of

ferromagnetic and paramagnetic component is reported due to some defect states present

in the doped ZnO systems. Sharma et al. [10] observed that there is coexistence of

ferromagnetic contribution and paramagnetic contribution in Zn(Fe)O system and the

paramagnetic behavior increased with increasing doping concentrations. They have

showed it by the Mossbauer spectra for Zn(Fe)O samples recorded at room temperature

in order to probe local magnetic environment around the Fe sites and to determine the

oxidation state of the Fe in ZnO matrix. Each spectrum shows a paramagnetic doublet

with isomer shift (IS). This coexistence of ferromagnetic and paramagnetic components

in other DMS systems has also been established by different researchers [16-18].

In this chapter, a systematic study of carrier concentration dependent room

temperature ferromagnetism (RTFM) in pulsed laser deposited epitaxial thin films of iron

doped zinc oxide [Zn(Fe)O] and iron doped zinc oxide incorporated with 1% aluminium

[Zn(Fe,Al)O] has been presented. We have investigated explicitly the presence of

temperature dependent paramagnetic component and its effect on room temperature

ferromagnetism in Zn(Fe)O and Zn(Fe,Al)O dilute magnetic semiconductors (DMS)

using low temperature SQUID measurements. The structural and optical absorption

properties have also been studied for these DMS thin films. Moreover, the electrical and

magneto-electrical properties have also been investigated for those DMS thin films. The

Anomalous Hall Effect and magnetoresistance behavior of these DMS films have been

studied and the sp-d exchange behavior of such DMS systems in different temperatures

has been found.

3.2. Experimental procedures

3.2.1. Preparation of targets

Required amount of high purity ZnO, Fe2O3 powder has been well mixed with

hand grinder repeatedly and sintered at 450 ºC till the required phase appeared with 5%


Chapter 3

46

Fe. Required amount of high purity Al2O3 powder were mixed with the ZnO and Fe2O3

powder to dope 5% of Fe and 1% aluminum in the zinc oxide target. The well mixed

powder after repeated grinding is sintered at the same 450 ºC. Finally, the Zn0.95Fe0.05O

and Zn0.94Fe0.04Al0.01O powders were palletized and used as the target for pulsed laser

deposition. The same procedures have been followed to dope ZnO with 7 and 10% Fe

also.

3.2.2. Cleaning of substrates

The c-plane (0001) sapphire substrate has been cleaned repeatedly with De-

ionized water, Acetone and Propanol using ultrasonic vibrator. Each cleaning process has

been carrier out for 20 min.

3.2.3. Preparation of thin films

The films are grown on well cleaned c-plane (0001) sapphire substrate by pulsed

laser deposition technique at several substrate temperatures (300 to 600 ºC) and different

ambient oxygen pressures (from base pressure ~10-5 to 10-1 Torr). The optimized

deposition parameters used in obtaining the Zn(Fe,Al)O films having highest magnetic

moment at room temperature are - substrate temperature of 450 ºC and 10-5 Torr ambient

oxygen atmospheres for 30 min at a laser pulse frequency (repetition rate) of 10 Hz. The

XeCl (λ = 308 nm) pulsed mode excimer laser has been used at an average pulsed laser

energy of 150 mJ.

3.2.4. Characterization of thin films

The concentrations of Fe were estimated by EDAX. The structural studies have

been done by high resolution x-ray diffraction (XRD), cross-sectional High Resolution

Transmission Electron Microscope (HRTEM). Atomic Force Microscope (AFM) image

have been taken for estimate the surface roughness of the films. Room temperature XAS

spectra have been taken to estimate the valence state of Fe ion and the content of metallic

Fe in the films. The magnetic properties have been carried out using SQUID

measurements at room temperature and low temperatures down to 5 K. The thicknesses

of the films have been measured with Dektak profilometer and found to lie in the range of

150 – 450 nm. Hall Effect and magnetoresistance measurements at room temperature and


47

low temperature have been performed using the van-der Pauw four-probe configuration

with magnetic fields up to 8 T. The band gap and crystallinity have been estimated using

transmission curve of UV spectroscopy recorded at different temperatures.

3.3. Results and discussions

3.3.1 Chemical properties study

The concentrations of Fe were established by energy dispersive x-ray analysis

(EDAX). The actual values of the Fe concentration in different films have been listed in

the Table-3.1.

Table-3.1: The actual Fe concentration in the Zn(Fe)O and Zn(Fe,Al)O thin films

Sample Actual Fe concentration

Zn(Fe)O with 5% Fe 2.85%

Zn(Fe,Al)O with 5% Fe 3.28%





3.3.2. Structural properties

In Fig. 3.1(a) the recorded XRD patterns (normal θ - 2θ scan) using Cu-Kα

radiation (λ = 1.542 Å) show that the Zn(Fe)O and Zn(Fe,Al)O films with 5% Fe of

thickness 420 and 300 nm respectively, are perfectly epitaxial on (0001) c-plane sapphire

in the (0001) direction of the films. The observed XRD patterns also confirm that there is

no iron oxide or any other secondary impurity phases present in the films. The ionic radii

of aluminum (Al3+) is smaller compared to zinc (Zn2+) but ionic radii of iron (Fe2+) is

larger than Zn. So on doping with iron in Zn(Fe)O a stress will develop in ZnO. The

FWHM of (0002) film peak and c-axis parameter are obtained to be 0.47˚ and 5.253 Å,

respectively. On co-doping with Fe and Al in ZnO as Zn(Fe,Al)O, the FWHM of (0002)


Chapter 3

48

film peak is obtained as 0.39˚ and c-axis parameter so obtained is 5.218 Å. The FWHM

and c-axis parameter are found to reduce slightly on co-doping as stress developed due

to doping of Fe and strain developed due to doping with Al. This development of stress

and strain compels the reduction of FWHM and c-axis parameter of the Zn(FeAl)O film

compared to FWHM and c parameter of the Zn(Fe)O film with 5% Fe. The XRD pattern

of Zn(Fe)O and Zn(Fe,Al)O thin films with different doping concentrations have been

shown in Fig. 3.1(b) and Fig.3.1(c), respectively.

Fig. 3.1(d) shows the high resolution transmission electron microscope (HRTEM) image

revealing very sharp film-substrate interface. The d spacing of the film calculated from

the HRTEM image is 0.261 which well matches with the d spacing of (0002) plane of

ZnO. The micrograph taken in [10-10]s || [2-1-10]f zone axes clearly indicates the high

degree of texturing of this epitaxial film with the substrate having an atomically sharp

interface with no mixed layer near the interface.

40 50 60 70

ZnO

2θ (degree)

Zn(Fe)O 5% Fe

Inte

nsity

Zn(FeAl)O 5% Fe

(a)

40 50 60 70

2θ (degree)

Zn(Fe)O 10% Fe

Zn(Fe)O 7% Fe In

tens

ity

(b)Zn(Fe)O 5% Fe

40 50 60

Zn(Fe,Al)O 10% Fe

Zn(Fe,Al)O 7% Fe

Zn(Fe,Al)O 5% Fe

2θ (degree)

(c)

Inte

nsity

Fig. 3.1. (a) The XRD pattern of as grown ZnO, Fe- doped and Fe and Al – doped epitaxial films for 5% Fe. (b) The XRD pattern of Zn(Fe)O epitaxial films for different Fe concentrations. (c) The XRD pattern of Zn(Fe,Al)O epitaxial films for different Fe concentrations. (d) the cross-sectional HRTEM of the junction.

(d)


49

Figure 3.2 shows the room temperature XAS spectra for both the Fe and Fe with

Al doped epitaxial films along with some standards. The valence of Fe in both the DMS

films appears to be Fe2+ as the band edge positions are similar to FeO spectra. The small

pre-edge peak of the films is likely due to the less symmetric environment in the Zn site

compared to the octahedral coordination in FeO.

3.3.3. Surface morphology

7110 7120 7130 7140 7150 71600.0

0.5

1.0

1.5

Fe2O3 maghemite Zn(Fe)OZn(FeAl)O FeO Fe2O3 Hametite

Nor

mal

ized

Χμ

(E)

Energy (eV)

Fig.3.2. (a) Room temperature near edges XAS spectra for both Zn(Fe)O and Zn(Fe,Al)O epitaxial films compared to some standards (Fe2O3 – Hematite, Fe2O3 –Maghemite and FeO). The valence looks to be Fe2+ for both the ZnO films.

10.80.60.40.20

1.2

1

0.8

0.6

0.4

0.2

0

X[µm]

Y[µ

m]

39.13 nm

0.00 nm1.41.210.80.60.40.20

1.2

1

0.8

0.6

0.4

0.2

0

X[µm]

Y[µ

m]

54.99 nm

0.00 nm

1.61.41.210.80.60.40.20

1.6

1.4

1.2

1

0.8

0.6

0.4

0.2

0

X[µm]

Y[µ

m]

59.16 nm

0.00 nm

(a) (b)

(c) (d)

(d)

Fig.3.3. (a), (b), (c) The AFM image of Zn(Fe)O and Zn(Fe,Al)O thin films. (d) 3-d view of one of the films.


Chapter 3

50

The AFM image of the films recorded for 1 µm × 1 µm scan area have been

shown in Fig. 3.3 (a) to (c) for Zn(Fe)O and Zn(Fe,Al)O thin films. The measured r.m.s

roughnesses of the films have been obtained to be 1 to 2 nm. Fig. 3.3 (d) is the 3-d view

of one of the scanned AFM image of Zn(Fe)O film.

3.3.4. Optical properties

To obtain an idea of the band gap of the films and its temperature variation, the

temperature dependent absorption spectra in UV regime have been recorded. The UV

transmission spectrum of Zn(Fe)O and Zn(Fe,Al)O films with 5% Fe at room

temperature are shown in the inset of Fig. 3.4(a). Fig. 3.4(b) and (c) are the same plot for

7 and 10 % Fe concentrations, respectively. The transmittance spectra show the well

crystalline nature (sharp drop at the band edges) of all the films. The band gap is

determined to be about 3.20 eV for all the films at room temperature.

The temperature dependence of the direct band gap, determined from the

absorption edge, can be described well by Varshni’s equation,

βγ+

−=T

TETE gg

2

)0()( (3.1)

The temperature dependent band gap shown in Fig. 3.5(b) can be fitted by Varshni’s

equation with Eg(0) = 3.55 eV, γ = 2.41 meV/K, and β = 935 K. In Varshni’s equation, β

is physically associated with the Debye temperature of the crystal and γ associated with

Fig.3.4. Room temperature transmission spectra of UV absorption spectroscopy for Zn(Fe)O and Zn(Fe,Al)O films (a) 5% Fe, (b) 7% Fe and (c) 10% Fe. (d) Room temperature transmission spectra of UV absorption spectroscopy for Zn(Fe)O films with different Fe concentrations.

200 400 600 800 1000354045505560657075808590

Zn(Fe)O Zn(Fe,Al)O

% T

rans

mitt

ance

Wavelength (nm)

5%

(a)

300 600 90030

40

50

60

70

80

90

(b)

7%

Zn(Fe)O Zn(Fe,Al)O

% T

rans

mitt

ance

Wavelength (nm)300 600 900

30

40

50

60

70

80

90

(c)

10%

Zn(Fe)O Zn(Fe,Al)O

% T

rans

mitt

ance

Wavelength (nm)


51

thermal expansion. The Debye temperature obtained for Fe, Al co-doped ZnO is

comparable to bulk undoped ZnO (920 K) [19].

3.3.5. Magnetic properties

The ferromagnetic M(H) behavior of both the Zn(Fe)O and Zn(Fe,Al)O films

grown on sapphire substrate at optimum deposition condition have been characterized

using a SQUID magnetometer in the magnetic field range of 0 - ± 5 T. The SQUID

measurements have been carried out down to 5 K operating temperature.

3.3.5.1. Room temperature magnetic properties

The ferromagnetic M(H) behavior at room temperature of both the Zn(Fe)O films

for different Fe doping concentrations have been shown in Fig. 3.6. The diamagnetic

contributions of sapphire substrate have been subtracted carefully at each magnetic field

from the net magnetization [uncorrected raw data (shown in the inset of Fig. 3.6.)] to

estimate the actual ferromagnetic contribution of each ferromagnetic film at 300 K. After

correcting the substrate contributions in SQUID raw data a ferromagnetic hysteretic

M(H) behavior at room temperature is observed for both the films. The coercive field and

saturation magnetization for the Zn(Fe)O samples have been listed in Table-3.2.

2.0 2.5 3.0 3.5 4.00

200

400

600 10 K 50 K 100 K 150 K 200 K 250 K 300 K 350 K 400 K

(αhν

)2

Energy (eV)

(a)

0 100 200 300 400

3.25

3.30

3.35

3.40

3.45

3.50

3.55

Ban

d ga

p (e

V)

Temperature (K)

(b)

Fig. 3.5. (a) (αhν)2 vs. energy plot of Zn(Fe,Al)O films at different temperatures. (b) Temperature dependent band gap plot for Zn(Fe,Al)O sample.


Chapter 3

52

Table-3.2: List of saturation magnetization and coercive field of Zn(Fe)O thin films with different Fe concentrations

Fe concentrations in

Zn(Fe)O thin films

Saturation magnetization

(μB/Fe2+)

Coercive field (Oe)

5% 0.18 135

7% 0.04 93

10% 0.02 87

If all the Fe2+ spins are aligned in high moment state one should expect to attain their full

saturation moment value of 4 μB/Fe2+ for all the films. The rather small value for

saturation magnetization suggests that only a small portion of Fe spins are probably

coupled ferromagnetically and that a significant paramagnetic and antiferromagnetic

fraction of Fe remains: this is seen explicitly from the observation that the diamagnetic

term subtracted from the total signal differs from what should be expected for a sapphire

substrate. Moreover, it is also well known fact that the oxygen vacancies produce shallow

donor states (defect states) while the zinc vacancies produce shallow acceptor states in

Fig.3.6. Room temperature ferromagnetic M(H) hysteresis loops for Zn0(Fe)O epitaxial films with 5, 7 and 10% Fe concentrations. Lower inset shows the same M(H) plot at 300 K in low field regime. Upper inset shows the uncorrected (from substrate contribution) SQUID raw data

-10000 -5000 0 5000 10000-0.15

-0.10

-0.05

0.00

0.05

0.10

0.15

10%

7%

5%

M (μ

B/F

e2+ )

H (Oe)

-400 -200 0 200 400

-0.04

-0.02

0.00

0.02

0.04

M (μ

B/F

e2+ )

H (Oe)

-10000 -5000 0 5000 10000-2

-1

0

1

2

M X

10-4

(em

u)

H (Oe)

Zn(Fe)O on sapphire


53

ZnO thin films. These states are delocalized due to the hybridization with the Fe d states.

The enhanced electron due to oxygen vacancy accommodates in the minority spin

channel (let say spin up), so minority spin channels become fully occupied e↑ level and

singly occupied t2↑ levels. On the other hand, Zn vacancies introduces holes into the

system, resulting in a completely empty minority spin channel and one hole in the

majority spin channel [18]. If the d orbital is partly occupied, the electrons in that orbital

then hop to the neighboring d orbital, and make the neighboring Fe atoms in parallel spin

configuration. It causes ferromagnetism in the film. On the other hand, if the d cell is

completely occupied, energy starts reduced via hoping process which causes

antiferromagnetic ordering in the system. This may be the cause of the loss of huge

amount of ferromagnetic moment from its expected value in our Fe doped ZnO films.

The low temperature M-H loop establishes the fact that the films may contain high order

of antiferromagnetic and paramagnetic moments. The decrease of ferromagnetic moment

with increasing concentration of iron may be due to the increase of antiferromagnetic

coupling between Fe pairs in the matrix. With increase in the Fe doping in ZnO, the

average distance between adjacent Fe2+ ions reduces. As the antiferromagnetic energy is

less than ferromagnetic energy, the antiferromagnetic coupling between Fe2+⎯Fe2+ ions

dominates at higher Fe concentrations and act as a ferromagnetic moment killer reducing

average magnetic moment per Fe ion. Similar results are obtained for Mn doped and Ni

doped ZnO films [20,21].

3.3.5.2. Low temperature magnetic properties

The temperature dependent M(H) raw SQUID data has been shown in Fig. 3.7(a).

The resulting graph of film magnetization versus magnetic field consists of contributions

from both the sample and the substrate. Measurements of uncoated sapphire substrates

show a large diamagnetic contribution [inset of Fig. 3.7(a)]. The raw SQUID data from a

Zn(Fe)O on sapphire shows two clear contributions [shown in Fig. 3.7(a)]. There is a

prominent low field ferromagnetic contribution from the epitaxial film. A high field

linear dependency with negative slope of diamagnetic contribution from substrate is also

observed. There can be a paramagnetic contribution from unreacted components in the

films which is also linear in this range. To separate out the ferromagnetic contribution


Chapter 3

54

from the rest, we use the fact that it saturates at higher magnetic fields. The

magnetization is linear above a certain field in the high field region and a straight line

with negative slope can be best fitted. At each temperature the slope of this line has been

calculated from the linear fit and subtracted from each point to leave the ferromagnetic

contribution. Thus obtained the ferromagnetic moment at different temperatures has been

shown in Fig. 3.7(b). The top-left inset of Fig. 3.7(b) shows the clear hysteresis loop at

temperature 5 and 300 K confirms the ferromagnetic behavior of the system. The bottom-

right inset of Fig. 3.7(b) shows the temperature dependent coercive field of Zn(Fe)O

samples.

The change of slope of the raw SQUID [M(H)] data in the high field region with

temperature as shown in Fig. 3.7(a) implies that the films contain not only the

diamagnetic substrate component (which is temperature independent) but also a huge

paramagnetic component in it. To separate out the susceptibility of paramagnetic

component at each temperature we consider that the high field slope of the linear fit

( paradia+χ ) contains both the diamagnetic and paramagnetic components. So,

paradiaparadia χχχ +=+ (3.2)

-40000-20000 0 20000 40000

-6

-4

-2

0

2

4

6

5 K 50 K 100 K 300 K

Mom

ent (

x10-4

emu)

Magnetic field (Oe)

-4 -2 0 2 4

-8

-4

0

4

8

Mom

ent(x

10-4

emu)

Magnetic field (x 104 Oe)

-60000 -40000 -20000 0 20000 40000 60000-1.5

-1.0

-0.5

0.0

0.5

1.0

1.5

5 K 50 K 100 K 300 K

Magnetic field (Oe)

Ferr

omag

netic

mom

ent

(x 1

0-4em

u)

-2000 -1000 0 1000 2000

-0.2

0.0

0.2

5 K 300 K

Magnetic field (Oe)

0 50 100 150 200 250 3000

200

400

600

800

HC (O

e)

Temperature (K)

Fig. 3.7. (a) Raw SQUID data at different temperature of Zn(Fe)O epitaxial thin film grown on c-sapphire substrate. Inset shows the SQUID data of blank substrate. (b) Ferromagnetism in different temperature. The top left inset is the low field hysteresis loop of Zn(Fe)O at temperature 5 K and 300 K. Right bottom inset shows the temperature dependent coercive field.

(a) (b)


55

where the slopes from raw SQUID data at a particular temperature in the very high field

regime give paradia+χ . diaχ and paraχ are the susceptibilities of diamagnetic component

from the substrate and film paramagnetic components which is most likely present from

unreacted component in the films. Now, diaχ is a temperature independent parameter and

paraχ depends on temperature with simple Curie law,

TC

para =χ (3.3)

where C is Curie constant and T is temperature in absolute scale (K). Combining Eq.

(3.1) and (3.2) we can write,

diaparadia TCT χχ +=+ (3.4)

The plot of paradiaT +χ with temperature [shown in Fig. 3.8(a)] gives a straight line plot

whose slope and intercept to y-axis gives the diaχ and C, respectively. The constant C

and diaχ has been found to be 8.38×10-7 emu/Oe/K and -2.74×10-8 emu/Oe, respectively.

We have also extracted the ferro, para and diamagnetic components of the same Zn(Fe)O

film by employing the M-H curve of blank (uncoated) c-sapphire substrate as shown in

inset of Fig. 3.7(a). The temperature dependent paramagnetic moment has been shown in

Fig. 3.8(b). The comparative studies of diamagnetic, paramagnetic and ferromagnetic

moment at 5 K and 300 K have been shown in upper left and lower right insets of Fig.

3.8(b). The magnetization in ferromagnetic materials is generally expressed as,

)(xJBNgM JBμ= where, Brillouin function ( ) ⎟⎠⎞

⎜⎝⎛−

++= x

JJx

JJ

JJxBJ 2

1coth21)

212coth(

212 and

TkJBgx

B

Bμ= . N is the number of atoms per unit volume, g is the g-factor and μB is the Bohr

magneton. J is described as the total angular momentum quantum number of the

microscopic magnetic moments of the material The saturating nature of paramagnetic

moment at higher fields measured at 5 K fits well with the expression of the

magnetization containing Brillouin function term keeping g =2 as shown in Fig. 3.8 (c).

The rest constant part (NgμB) in the expression of the magnetization has been described

as constant A.


Chapter 3

56

Table-3.3: Obtained fit parameters from the Brillouin function fit keeping g =2 at different temperatures.

Temperature (K) J A R2(Goodness of fit)

5 1.21 ± 0.04 0.0011 ± 1 х 10-5 0.99975

50 1.19 ± 0.13 0.0008 ± 5 х 10-5 0.99978

100 0.98 ± 0.11 0.0007 ± 4 х 10-5 0.99986

300 0.991 ± 0.12 0.00034 ± 2 х 10-5 0.99979

We get J~1 with excellent goodness of fitting (R2= 0.99975). We have also fitted the

paramagnetic moments with the Brillouin function keeping g = 2 for different

temperatures also. The fitted parameters for different temperatures (5 K, 50 K, 100 K and

0 50 100 150 200 250 300-8

-6

-4

-2

0

Tc d

ia+p

ara

(x 1

0-6)

Temperature (K)

(a)

-10000 -5000 0 5000 10000-8-6-4-202468

(b)

Para

mag

netic

m

omen

t (x1

0-5 e

mu)

Magnetic field (Oe)

5K 50K 100K 300K -8000 -4000 0 4000 8000

-20

-10

0

10

20

Para

Magnetic field (Oe)

Mon

ent (

x 10

-5em

u)

Dia

Ferro

300 K

-70000 -35000 0 35000 70000

-8

-4

0

4

8

DiaFerro

Mom

ent (

x10-4

em

u )

Magnetic field (Oe)

T=5 K

Para

-0.8 -0.4 0.0 0.4 0.8

-8

-4

0

4

8

(c)

Para

mag

netic

mom

ent (

x 10

-4em

u)

μB

/kBT

5 K

0 50 100 150 200 250 300

0.30.60.91.2

02468

II

Temperature (K)

Ms (

x10-4

emu)

(d)

χ par

a (x

10-9

em

u/O

e)

I

Fig. 3.8. (a) The plot of TT paradia −+χ . (b) Extracted paramagnetic moment at different

temperatures. Insets are the comparative study of diamagnetic, paramagnetic and ferromagnetic moments at 5 K (upper left inset) and 300 K (lower left inset) of the same film. (c) The paramagnetic moment at 5 K fitted with Brillouin function keeping g=2. (d) Paramagnetic susceptibility (I) and saturation magnetization (II) vs. temperature plots of Zn(Fe)O film.


57

300 K) have been summarized in Table-3.3. The J values obtained from the fit for all

temperatures are found to be ~ 1.0. From the excellent fit (R2 ≈ 1) as shown in Table-3.3,

it concludes that the paramagnetic moments at different temperatures can be fitted well

with a single Brillouin function with a constant value of J (~1). The small value of J

indicates that it does not approach towards classical limits (Langevin function). The

parameters extracted from the Eq. (3.3) and experimental result has been shown in the

Table-3.4. The diamagnetic susceptibility contribution is almost comparable (Table-3.4)

but mismatch of the paramagnetic susceptibilities has been found in the system.

Table-3.4: Comparative study of paramagnetic susceptibilities at different temperatures

Temperature (K)

Paramagnetic susceptibility calculated using the

measurement of diamagnetic susceptibility of sapphire

substrate(-3.11×10-7 (emu/gm)/Oe)

Paramagnetic susceptibility calculated using the Eq. (3.4).

Calculated diamagnetic susceptibility (-8.32×10-7

(emu/gm)/Oe)

5 4.43×10-9 1.67×10-7

50 1.78×10-9 1.67×10-8

100 14.73×10-10 8.35×10-9

300 3.02×10-10 2.78×10-9

The mismatches of the paramagnetic susceptibilities of the films, estimated from Eq. 3.4

and experiment, imply that the simple Curie law cannot be applicable directly for the

DMS systems. The temperature dependent ferromagnetic saturation moment (Ms) and

paramagnetic susceptibility (χpara) has been shown in Fig 3.8 (d). Both ferromagnetic

saturation moment and paramagnetic susceptibility decrease with temperature and they

do not follow standard Curie law (χ ∝ 1/T). A sharp exponential rise of both

ferromagnetic moment and paramagnetic susceptibility in the low temperature regime

reveals more like the insulating type DMS behavior of the film [8,22].

This huge paramagnetic moment may also appear in the films by defects. As

discussed earlier it is well known fact that the oxygen vacancies produce shallow donor

states (defect states) while the zinc vacancies produce shallow acceptor states in ZnO thin

films. These states are delocalized due to the hybridization with the Fe d-states. The


Chapter 3

58

enhanced electron due to oxygen vacancy accommodates in the minority spin channel

(spin up) of Fe, so minority spin channels become fully occupied eg↑ level and singly

occupied t2g↑ levels. On the other hand, Zn vacancies introduce holes into the system,

resulting in a completely empty minority spin channel and one hole in the majority spin

channel. If the d orbital is partly occupied, the electrons in that orbital then hop to the

neighboring d orbital, and make the neighboring Fe atoms in parallel spin configuration.

It causes ferromagnetism in the film. On the other hand, if the d cell is completely

occupied, energy starts reducing via hoping process which causes antiferromagnetic

ordering in the system.

From the temperature dependent M-H loop in Fig. 3.7(b), it is found that the

magnetization saturates at higher field at lower temperatures compared to same at room

temperature. According to the conventional theory of magnetization when thermal energy

(kBT) is less, the magnetization should saturate at lower fields. But the situation is

completely different in this case. This can be explained by the properties of pinning-type

magnets [23]. In pinning type magnet, the Bloch walls cannot travel freely throughout the

whole grain because of magnetic inhomogeneities present in the grains. These magnetic

inhomogeneities act as pinning centers for the domain walls motion. Apart from the

change in magnetization associated with some wall bending, this pinning will prevent

further magnetization. Wall displacement (other than bending) can occur only when the

force exerted on the wall becomes sufficiently strong. When the strength of the external

field exceeds the pinning field strength then only the saturation magnetization occurs. As

being very low thermal energy in low temperatures the saturation magnetic fields are

higher than the room temperature. So the M-H curve at 300 K saturates earlier than the

curves recorded at lower temperatures.

The isothermal M2 vs H/M plots (Arrott-Belov plot) have been shown in Fig. 3.9

(a) to confirm the presence of intrinsic spontaneous magnetization [M1/β vs. (H/M)1/γ

isothermal plots where β = 0.5 and γ=1 in the mean-field limit]. Each plot clearly shows

the positive intercept at y-axis for all the temperatures confirming the presence of

ferromagnetic spontaneous magnetization in the film for all the temperatures up to 300 K.


59

The temperature dependent spontaneous magnetization obtained from Arrott-Belov plots

has been shown in Fig. 3.9 (b). The temperature dependent spontaneous magnetization

curve follows the same sharp rise at low temperature behavior as saturation

magnetization follows with temperature [shown in Fig. 3.8 (d)]. The temperature

dependent behavior of spontaneous magnetization has been found to be similar to an

insulating type DMS material.

3.3.5.3. Carrier dependent ferromagnetism properties

Incorporating 1% of Al in the Zn(Fe)O matrix the room temperature

ferromagnetic component increases almost four times at room temperature as shown in

Fig 3.10(a). However, the understanding of the effect of the additional dopants on the

enhancement of ferromagnetism in the DMSs is still an issue of debate. In some literature

it has been reported that the additional carrier causes enhancement of ferromagnetism

[14]. Li et al. [24] suggest that introducing of Al in Mn doped ZnO results not the

increase of carrier concentration but it could break the metastable structures that formed

in Mn doped ZnO, leading to spinodal decomposition and lead to Mn rich regions.

0 50 100 150 200 250 3000.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

(b)

spon

tane

ous m

agne

tizat

ion

(em

u/g)

Temperature (K)

0.0 4.0x104 8.0x104 1.2x105 1.6x105

10-1

100

300 K

100 K50 K

M2

(em

u2 g-2)

H/M (Oe-g/emu)

Arrott-Belov plot5 K

(a)

Fig. 3.9. (a) Arrott-Belov plots for of Zn(Fe)O epitaxial thin film at different temperatures. (b) Temperature dependent spontaneous magnetization of Zn(Fe)O sample.


Chapter 3

60

The Mn rich region could be the cause of enhancement of ferromagnetic ordering. In our

case, we observe that the (Fe, Al) doped ZnO films show the higher carrier concentration

( ~ 8.02 × 1026 m-3) than the Fe doped one ( ~ 3.34 × 1026 m-3). The carrier concentrations

of all our DMS films were estimated from the magnetic field dependent Hall voltage

measurements at room temperature. As discussed in XRD results, it has been observed

that the introduction of Al in Fe doped ZnO releases the strain or lattice distortion in our

films. Al in Fe doped ZnO results in increase in its carrier density and also break the

metastable structures that formed in Fe doped ZnO, leading to spinodal decomposition,

which may cause the Fe rich regions. These regions may also be the cause of

enhancement of ferromagnetism due to incorporation of Al in iron doped ZnO films. We

0.0 0.2 0.4 0.6 0.8 1.0 1.2-0.20.00.20.40.60.81.01.21.4

(c)

M (e

mu/

gm)

nc/ni 0.0 0.2 0.4 0.6 0.8 1.0 1.2

0

2

4

6

8

10

(d)χ pa

ra (x

10-9

emu.

gm-1

Oe-1

)

nc/ni

-10000 -5000 0 5000 10000-0.4-0.3-0.2-0.10.00.10.20.30.4

Zn(Fe)O

Zn(FeAl)O

H (Oe)

M (μ

B/F

e2+)

(a)

-10000 -5000 0 5000 10000-0.6

-0.4

-0.2

0.0

0.2

0.4

0.6

(b)

Zn(FeAl)O

Zn(Fe)O

Para

mag

neic

mom

ent

(em

u/gm

)

Magnetic field (Oe)

Fig. 3.10. (a) Room temperature ferromagnetic moment M-H loop of Zn(Fe)O and Zn(Fe,Al)O films, (b) room temperature magnetic field dependent paramagnetic moment contribution of those films. (c) and (d) are carrier dependent (nc/ni) room temperature ferromagnetic moment and paramagnetic susceptibility, respectively.


61

have also studied the other Al incorporated films which have much higher carrier

concentrations (listed in Table-3.5).

Table 3.5. Comparative FWHM and carrier concentrations with magnetic moment

Samples Carrier Concentrations

nc/ni FWHM Magnetic moment

Zn(Fe)O 5% Fe 1.61× 1020 0.076 0.472˚ 0.04000

Zn(Fe)O 5% Fe 2.08× 1020 0.099 0.471˚ 0.03896

Zn(Fe)O 5% Fe 2.93× 1020 0.139 0.471˚ 0.08980

Zn(Fe)O 5% Fe 3.34× 1020 0.158 0.472˚ 0.08560

Zn(Fe,Al)O 5% Fe 3.09× 1020 0.147 0.394˚ 0.00000

Zn(Fe,Al)O 5% Fe 7.02× 1020 0.333 0.395˚ 0.38000

Zn(Fe,Al)O 5% Fe 7.64× 1020 0.362 0.394˚ 0.20500

Zn(Fe,Al)O 5% Fe 1.08× 1020 0.513 0.393˚ 0.04686

Zn(Fe,Al)O 5% Fe 1.25× 1020 0.593 0.394˚ 0.01976

Zn(Fe,Al)O 5% Fe 2.22× 1020 1.051 0.394˚ 0

The XRD study of those films also shows the strain relaxation which should cause the

spinodal decomposition, and hence should enhance the magnetic moment. But, those

films with higher carrier concentrations show lower moment which contradict the

previously discussed ‘spinodal decomposition’ theory [24]. Hence, the observation of

room temperature ferromagnetic behavior in our DMS films probably can be best

explained through the standard theory of carrier induced ferromagnetism. Carrying out

the low temperature SQUID measurements and separating out the diamagnetic,

paramagnetic, ferromagnetic components as already discussed above, it is clearly seen

that the ferromagnetic exchange interaction (magnetization) increases with carrier

concentration and reaches to a maximum beyond which it falls drastically at higher

carrier concentrations as shown in Fig. 3.10(c). Standard theory of DMS predicts that the

exchange is maximized when the nc/ni ratio is in between 0.3 to 0.5 [14]. The reduction of

the moment with increasing carrier density ratio of the films occurs when nc/ni > 0.4. In

addition, from the low temperature SQUID analysis the extracted paramagnetic


Chapter 3

62

susceptibility with carrier concentration shows exactly opposite nature as shown in Fig.

3.10(d).

The defect such as anion vacancy plays an important role in ferromagnetism in

ZnO as this defect forms shallow donor and provides n-type conduction. Zn interstitial is

also cause the shallow donor level. The cation vacancy can introduce an electronic defect

on neighboring oxygen and creates an O-2p5 anion. The oxygen vacancies are strongly

correlated, and they can form extended molecular orbit around the defect site, which

couples ferromagnetically. The proposed electronic structure with impurity band splits at

higher Curie temperature has been proposed by Coey et al. [25] as shown in Fig. 3.11(a).

Doping with Fe2+ splits d level of shallow impurity band which moves down to the 2p

band of oxygen. Hence, two region forms: one near the beginning of the series where 3d↑

state cross Fermi level in the impurity band, and one towards the end where 3d↓ state

d4↑ d4↑ d4↑ d4↑

d4↑ d4↑ d4↑ d4↑

d4↑ d4↑ d4↑ d4↑

d4↑ d4↑ d4↑ d4↑

d5↑ d5↑ d5↑ d5↑

d5↑ d5↑ d5↑ d5↑

d5↑ d5↑ d5↑ d5↑

d5↑ d5↑ d5↑ d5↑

d5↑ d10 d10 d5↑

d10 d10 d10 d10

d10 d10 d10 d10

d5↑ d10 d10 d5↑

(a) (b)

(c) (d)

e- e-

Fig.3.11.(a) Electronic band structure for high Tc oxides proposed by Coey et al. [25], (b) The magnetic region in Zn(Fe)O sample. (c) After introducing excess carrier to the system within the limit. The excess electrons sit on the unfilled d state and enhance the up spin population of the polarons. (d) Condition when the carrier density is much higher. The carrier density is higher (cross the limit of nc/ni) the d cell starts to be filled by electrons and decreases up spin population. It causes decrease of magnetic moment. When the electron concentration is very high the d cells of maximum polarons become completely filled and the polaron-polaron distance increases which causes reduction of magnetic moment.


63

cross the Fermi level. If the impurity band is sufficiently narrow, the donor electrons

become localized and interact with core spin. Hence, if the localized electron interacts

with many magnetic cations it shows completely spin polarization. As the net moment of

iron doped samples are much less as expected (4µB), one can concludes that the localized

electrons are not completely able to interact with all magnetic cations. We also observed

that with increasing carrier concentration the magnetization increases. But additional

doping of electron kills the magnetic moment beyond a certain value of carrier

concentration (nc/ni ~ 0.4). In Fig. 3.11(b), (c) and (d) we have discussed this mechanism.

The iron doped ZnO films contains magnetic cation with partially filled d4 states (Fig.

3.11(b)). The magnetization occurs from the interaction of those cations with localized

electrons. If free carrier increased in the limit nc/ni < 0.4, the additional electrons transfers

to the unfilled d cell and increase net magnetic moment (Fig 3.11(c)). If the carrier

concentrations increases further beyond the limit, the excess carriers enter to d states and

decreases half field d shell. Net magnetization starts decreasing. At very high carrier

concentration d shells become completely filled and net Fe2+ ions available in the system

is reduced. So, the distance between nearest magnetically active ions become so large

that cannot mediate long range magnetic ordering. The films become eventually

paramagnetic in nature.

3.3.6. Electrical properties

The presence of magnetic ions such as 3d transition metal (TM) ions in these

materials leads to an exchange interaction between traveling sp band electrons or holes

and the d electron spins localized at the magnetic ions, resulting in versatile magnetic-

field-induced functionalities [26]. DMS based materials demonstrated several spin related

phenomena such as a spin-polarized transport and luminescence attributable to a sp-d

exchange interaction [27]. In the recent years, ZnO based dilute magnetic semiconductors

have attracts a lot because of huge controversy between the DMS researchers. So, ZnO

based DMS study has a great opening for the researchers.

The anomalous Hall Effect (AHE) has been recognized as a powerful technique

for demonstrating the ferromagnetic ordering to be intrinsic due to spin-polarized

carriers, which mediate ferromagnetic exchange interaction with localized magnetic


Chapter 3

64

moments. AHE is well-known to be caused by the emergence of voltage transverse to

both an applied current and an external magnetic field proportional to magnetization [28].

The anomalous Hall Effect is attributed to asymmetric scattering involving in spin orbit

interaction between conduction electrons and the magnetic moments [29,30]. So this

analysis is widely used to study the magnetic behavior of dilute magnetic semiconductors

[31,32]. However, the exact mechanism is not clear so far. The questions arise whether

the mechanism is intrinsic or extrinsic. Karplus and Luttinger [33] have proposed the

phenomenon of Anomalous Hall Effect is intrinsic. On the other hand, some theories

attribute the impurity scattering modified by spin orbit interaction namely the skew

scattering [34] and the side jump mechanism [35]. This extrinsic mechanism is rather

complicated and depends on the impurities and band structures of the DMS materials

[36].

One of the characteristics of features of magnetic semiconductors is the sp-d

exchange between sp band of the semiconductor and the localized d electrons associated

with magnetic ions. Magneto-transport measurements in the DMS systems are

extensively used for studying the sp-d interactions of ferromagnetic semiconductors [37,

38]. It is well known fact that the magneto-transport is completely dependent of free

carrier concentration of n-type ferromagnetic semiconductors like ZnO [39-41]. The

change in carrier concentrations and as well as temperature the characterization of wave

function changes from delocalized to localized states which are responsible for change in

magnetoresistances (MR) in the DMS systems [42]. Several different mechanisms have

been proposed for the spin-dependent MR effect [43,44]. One of the mechanisms is the

formation of magnetic polarons [45]. The concept of magnetic polarons can be

understood as the variations in the magnetization created by those doped magnetic ions

around the localized carriers [46]. If the electrons want to travel through the lattice with

low resistance, it must carry the same spin polarization as the surrounding magnetization.

Clearly, the transport behavior is influenced by the magnetization condition of the films.

3.3.6.1. Electrical transport properties

The temperature dependent resistivity of the films has been shown in Fig. 3.12 for the

temperature range of 1.5 to 300 K. The resistivity drops with temperature for all the


65

films. Figure 3.12(a) and Fig. 3.12(b) are the temperature dependent resistivity plot of the

Zn(Fe)O and Zn(Fe,Al)O films with different iron doping.

The conduction mechanism in these DMS films can be best explained by the combination

of three types of conduction models

(i) variable range hopping (VHR),⎥⎥

⎦

⎤

⎢⎢

⎣

⎡

⎟⎟⎠

⎞⎜⎜⎝

⎛−=

4/10

0 expT

TVRH

VRHVRH σσ (3.5),

which describes carrier hopping in localized states.

(ii) Efro’s variable range hopping (EVRH), ⎥⎥

⎦

⎤

⎢⎢

⎣

⎡

⎟⎟⎠

⎞⎜⎜⎝

⎛−=

2/10

0 expT

TEVRH

EVRHEVRH σσ

(3.6) due to the electron-electron Coulomb interaction at lower

temperature range.

0 50 100 150 200 250 3002.0x10-4

3.0x10-4

4.0x10-4

5.0x10-4

6.0x10-4

7.0x10-4

8.0x10-4

10% Fe

7% Fe

ρ (Ω

m)

Temperature (T)

Zn(Fe)O

5% Fe

0 50 100 150 200 250 3006.0x10-5

8.0x10-5

1.0x10-4

1.2x10-4

1.4x10-4

1.6x10-4

1.8x10-4

2.0x10-4

2.2x10-4

2.4x10-4

10%Fe7%Fe

ρ (Ω

m)

Temperature (K)

Zn(FeAl)O

5% Fe

4 5 6 7 8 9 10-10.0

-9.6

-9.2

-8.8

-8.4

-8.0

-7.6

-7.2

Zn(FeAl)O

Zn(Fe)O

ln (ρ

)

1000/T (K-1)

0.2 0.3 0.4 0.5 0.6 0.7 0.8-10.8-10.4-10.0-9.6-9.2-8.8-8.4-8.0-7.6-7.2

Zn(FeAl)O

ln (ρ

)

T-1/4 (K-1/4

)

Zn(Fe)O

0.120 0.144 0.168 0.192-9.6-9.2-8.8-8.4-8.0-7.6-7.2

Zn(FeAl)O

Zn(Fe)O

T-1/2(K-1/2)

ln (ρ

)

Fig. 3.12. (a) Temperature dependent resistivity plot of Zn(Fe)O thin films with 5, 7 and 10% Fe, (b) The same plot for the 1% Al incorporated samples, (c) ln(ρ) - T-1/4 plot of Zn(Fe)O and Zn(Fe,Al)O thin films with 5% Fe to test the VRH mechanism. Inset is the ln(ρ) - T-1/2 plot of those same films to test the EVRH mechanism and (d) ln(ρ) -1000/T plot of those samples which satisfies the thermal excitation process of transport.

(a)

(b)

(c) (d)


Chapter 3

66

(iii) Thermal excitation model, ⎥⎦

⎤⎢⎣

⎡⎟⎟⎠

⎞⎜⎜⎝

⎛−=

kTE A

th thexp0σσ (3.7),

which describes carriers that have been thermally excited from

localized sate to conduction band.

VRH0σ , VRH

T0 , EVRH0σ ,

EVRHT0 and th0σ are constants. EA is the thermal activation energy. The

total resistivity can be expressed as,

( ) 1−++= thEVRHVRH σσσρ (3.8)

The parameters evaluated from the best non-linear χ2 fitting method have been listed in

the Table-3.6.

Table-3.6. Evaluated fit parameters of ρ-T behavior of the Zn(Fe)O and Zn(Fe,Al)O

samples using Eq. (3.8).

Samples σ0VRH (Ωm)-1

T0VRH (K)

σ0EVRH (Ωm)-1

T0VRH (K)

σ0Th (Ωm)-1

EA (eV)

Adj. R2

Zn(Fe)O 5% Fe 42423 32591 1572 0.032 19426 0.08256 0.99986 7% Fe 12707 15105 1619 0.015 21912 0.0903 0.99927 10% Fe 2968 2474 1429 0.045 4673 0.05779 0.99993

Zn(Fe,Al)O 5% Fe 2650208 367993 6035 0.29 1345018 0.31743 0.99895 7% Fe 830367 203006 6295 0.031 2075335 0.21689 0.99972 10% Fe 13230 2474 6368 0.045 20828 0.05779 0.99993

To investigate the temperature dependent transport process in the films we plot

ln(ρ) with T-1/4, T-1/2, and 1000/T in different temperature ranges. Figure 3.12(c) shows

the ln(ρ) vs T-1/4 plot of Zn(Fe)O film doped with 5% Fe to find the cross over

temperature (T*) and T0. The resistivity rises at low temperature and passes the VRH test

(T*<T0) at very low temperature range (1.6 to 4 K). Inset of Fig. 3.12(c) shows the ln(ρ)

plot with T-1/2 of the same film which gives a straight line in the temperature range 25 to

150 K, implies the Efros’s variable range hopping due to the electron-electron Coulomb

interaction at low temperature range [50]. In between 4 to 25 K the VRH is classed as

intermediate [14]. ln(ρ) with 1000/T plot of the Zn(Fe)O film has been shown in Fig.

3.12(d) for the temperature range 150 to 300 K. It suggests that the conduction in these


67

thin films is due to the thermally assisted tunneling of the charge carriers through the

grain boundary barrier and transition from donor level to conduction band. The transport

mechanism at different temperature ranges has been tabulated in Table-3.7.

Table-3.7. The summarized transport mechanism in different temperature ranges

Sample Temperature Range

Transport mechanism

Best fit R2 value

Zn(Fe)O 1.6-5 K Variable range hopping [T* < T0]

VRH test passed

5 -40 K Intermediate[T* >T0] VRH test failed 40-150 K Efros’s VRH

[ρ = ρ0 exp (T0/T)1/2 ] 0.970

150 – 300 K Thermal excitation

[ρ ~1000/T] 0.0987

Zn(Fe,Al)O 1.6-5 K Variable range hopping [T* < T0]

VRH test passed

5 -40 K Intermediate [T* >T0] VRH test failed 40-150 K Efros’s VRH

[ρ = ρ0 exp (T0/T)1/2 ] 0.988

150 -300 K Thermal excitation [ρ ~1000/T]

0.0992

3.3.6.2. Hall Effect study

The hall resistivity with magnetic field plot of Zn(Fe,Al)O film with 5% Fe at

different isothermal temperatures has been shown in Fig. 3.13(a). Fig. 3.13(b) and Fig.

3.13(c) are the same plot of the Zn(Fe)O and Zn(Fe,Al)O, respectively for different iron

concentrations at 2 K and 300 K (corresponding insets). All the films show the n-type

nature of the films. Saturating nature of the Hall resistivity at higher magnetic field range

in Fig. 3.13(a) confirms the anomalous Hall Effect behavior of iron doped epitaxial ZnO

films. The anomalous Hall Effect can be described by,

MRBR sH 00 μρ += (3.9)


Chapter 3

68

Where, Hρ , R0 and Rs are the Hall resistivity, ordinary Hall co-efficient and

anomalous Hall co-efficient, respectively. M is the magnetization of the film. The plot of

ordinary Hall co-efficient with temperature of all the films evaluated from temperature

dependent Hall measurements have been shown in Fig. 3.13(d) and its insets.

3.3.6.2.1. Ordinary Hall Effect

From the ordinary Hall co-efficient one can easily find the carrier concentrations and Hall

mobility of these DMS films. The calculated temperature dependent carrier concentration

of Zn(Fe)O and Zn(Fe,Al)O films for different iron concentrations have been plotted in

Fig. 3.14(a) and Fig. 3.14(b), respectively. The increase of carrier concentration with

0 2 4 6 80.0

0.5

1.0

1.5

2.0 2 K 5 K 10 K 50 K 100 K 150 K 200 K 250 K 300 K

ρ hall (

μΩ-m

)

B (T)

Zn(Fe,Al)O with 5% Fe

(a)

0 2 4 6 80.00.51.01.52.0

2.53.03.5

5% 7 % 10%

ρ hal

l (Ω

−m)

B (T)

T = 2 K

Zn(Fe)O

(b)

0 2 4 6 80.0

0.1

0.2

0.3

0.4

0.5

ρ ha

ll(μΩ

−m)

B (T)

T = 300 K

0 1 2 3 4 5 6 7 8 90.0

0.5

1.0

1.5

2.0

2.5(c) 5%

7% 10%

ρ hall (μ

Ω-m

)

B(T)

T =2 K

Zn(Fe,Al)O

0 2 4 6 80.0

0.1

0.2

0.3

ρ hall (μ

Ω−m

)

B (T)

T =300 K

0 50 100 150 200 250 3000.0

0.1

0.2

0.3

0.4

0.5

(d)

R0 (

m3 /C

)

Temperature (T)

Zn(Fe)O

0 50 100 150 200 250 3000.0

0.1

0.2

0.3

0.4

0.5 5% Fe 7% Fe 10% Fe

R0 (

μΩ−m

)

Temperature (T)

Zn(FeAl)O

Fig.3.13. (a) Magnetic field dependent Hall resistivity of Zn(Fe,Al)O thin film with 5% Fe measured at different isothermal temperatures. (b) and (c) are the magnetic field dependent Hall resistivity plots of Zn(Fe)O and Zn(Fe,Al)O thin films, respectively for different Fe concentrations measured at 2 K. Corresponding insets are the same plots measured at 300 K. (d) Temperature dependent ordinary Hall co-efficient of Zn(Fe)O thin films with different Fe concentration. Inset is the same plot for Zn(Fe,Al)O sample.


69

temperature shows semiconducting behavior of the films. To analyze the temperature

dependent carrier concentration behavior we have used the multi-donor charge balance

equation so that shallow electrically active defect concentrations and activation energies

could be extracted [47].

For non-degenerate n-type conduction, the charge balance equation adopts the form,

∑=

−

⎟⎟⎠

⎞⎜⎜⎝

⎛ Δ+

=+k

i

B

Di

c

Di

DiA

TkE

NnTg

NNn

12/3

exp1 (3.10)

where n is the carrier concentration, NA is the total acceptor concentration and NDi, ΔEDi

and gDi are the concentration, activation energy and donor degeneracy factor of the ith

donor Di, respectively. We consider gDi = 2, NC is the effective density of states in the

0 100 200 300 400 5000.00.20.40.60.81.01.21.41.61.8

FeAl5 FeAl7 FeAl10

n (x

1026

m-3

)

1000/T0 100 200 300 400 500

0.00.20.4

0.60.8

1.01.21.4

1000/T

n (x

1026

m-3

) Fe5 Fe7 Fe10

0 50 100 150 200 250 3001

2

3

4

5

6

7

Hal

l mob

ility

(x10

-4 m

2 /V-s

)

Temperature (K)

Zn(Fe)O

0 50 100 150 200 250 3004

8

12

16

20

24

28

H

all m

obili

ty (x

10-4

m2 /V

-s)

Temperature (K)

Zn(Fe,Al)O

(a) (b)

(c) (d)

Fig. 3.14. (a) and (b) are the carrier concentration evaluated from R0 plot with 1000/T for Zn(Fe)O and Zn(Fe,Al)O thin films with different Fe doping. (c) and (d) are the Hall mobility of those films.


Chapter 3

70

conduction band at 1 K given by 2/32* ))2/((2 hπBc kmN = , where m* is the density of states’

effective mass, kB is Boltzmann’s constant. To fit the temperature dependent carrier

concentration we have considered that the acceptor concentration is very less and ND ≈ n.

The fitted parameters for k = 2 have been shown in the Table-3.8.

Table-3.8. The evaluated fit parameters from temperature dependent carrier concentration using Eq. (3.10)

Sample ND1 (m-3) ND2 (m-3) ED1 (meV) ED2 (meV)

Zn(Fe)O 5% 1.07х1028 3.47 х1025 181.79 217.41

7% 9.83 х1027 1.89 х1025 133.61 156.15

10% 2.15 х1028 1.61 х1025 151.79 208.26

Zn(Fe,Al)O 5% 1.09 х1028 3.9 х1025 75.59 315.80

7% 2.16 х1028 3.36 х1025 108.92 333.59

10% 9.54 х1026 1.30 х1025 545.11 72.89

The temperature dependent Hall mobility has been plotted in Fig. 3.14(c) and Fig.

3.14(d) for Zn(Fe)O and Zn(Fe,Al)O, respectively. The mobility first decreases with

increasing temperature up to a certain temperature. After that it starts to increase up to a

certain temperature and then starts decreasing. The mobility of electrons in non-

degenerate single crystal ZnO is limited primarily by scattering due to ionized impurities,

deformation potential and piezoelectric acoustic phonons, and polar optic phonons. Grain

boundary scattering has to be considered as well for such films. The carrier mobility also

depends on free carrier concentration of the films. At very low temperature range the

electron-electron scattering dominates over other. With increasing temperature free

carrier concentration increases and causes a decrease of mobility. The electron-electron

collision do not affect the current density directly as they can not alter the total

momentum. They just randomize the carrier distribution and the momentum randomly

distributed to different velocity groups. The electron-electron collision gives rise to a net


71

transfer of momentum, and results greater rate of momentum transfer and lower the

mobility. The electron-electron scattering mobility is inversely proportional to free

electron concentration [ )(/1 Tnee ∝−μ ] [48]. In the comparatively higher temperature

range the electron-electron scattering does not dominates and the impurity scattering as

well as grain boundary scattering starts to dominate. In the moderate low temperature

region impurity scattering dominates. With increasing temperature the localized impurity

scattering start decrease which causes an increase of mobility. The impurity scattering is

given by [49],

1

2

222/31ln

−

⎥⎥⎦

⎤

⎢⎢⎣

⎡

+−⎟

⎟⎠

⎞⎜⎜⎝

⎛+=

BTNBT

NBT

NAT

IIIIμ (3.11)

Where, A and B are constants and NI is the ionized impurity scattering. The electron can

face some grain boundary scattering also. The mobility due to grain boundary scattering

is given by,

( )Tk BBB /exp0 φμμ −= (3.12)

where Tklcq B8/0 =μ . The l is the grain size and 2/1* )/8( mTkc B π= is the thermal

velocity. Bφ is the effective barrier height between two grains. At higher temperature

lattice scattering dominates which causes again decrease of mobility. The temperature

dependent mobility due to lattice scattering can be expressed as [50], α

μ ⎟⎠⎞

⎜⎝⎛=

TDCL (3.13)

Where C, D and α are constants. The total Hall mobility can be expressed using

Matthiessen’s rule,

LBIeeH μμμμμ11111

+++=−

(3.14)

3.3.6.2.2. Anomalous Hall Effect

Figure 3.15(a) shows the temperature dependence of anomalous Hall coefficient,

Rs of the Zn(Fe,Al)O film with 5% Fe, exhibiting decrease of Rs with increasing

temperature. The observation of anomalous Hall Effect in this DMS film also provides


Chapter 3

72

evidence of intrinsic carrier-mediated ferromagnetism [51]. The AHE observed in other

Zn(Fe)O and Zn(Fe,Al)O films with higher Fe doping is very weak compared to

Zn(Fe,Al)O film with 5% Fe. One of the possible reasons might be the relatively large R0

and low magnetic moments in other Zn(Fe)O and Zn(Fe,Al)O thin films due to smaller

carrier concentration [52]. The origin of the anomalous Hall Effect is believed to be the

spin-orbit interaction between the carrier angular momentum and the localized spin. Rs

has a power law relationship with the ohmic resistivity and is given by,

ns CR ρ= (3.15)

where, C is a constant. The exponent n=1 corresponds to the skew scattering and n=2

corresponds the quantum mechanical side jump scattering [53]. Considering the both

mechanism one can write [54],

2ρρ sjsks baR += (3.16)

where Masksk 0~ μφ represents the average deflection of a charge carrier at a scattering

center. The bsj is associated with a side-jump mechanism where the charge carrier’s

trajectory is displaced a fixed distance perpendicular to its original path at each scattering

centers.

1.0x10-4 1.5x10-4 2.0x10-4 2.5x10-4

0.00

0.01

0.02

0.03

R s=aρ

+bρ2

Rs (m

3 /C)

ρ (Ω-m)

Rs=bρ2

(b)

0 50 100 150 200 250 300

0.00

0.01

0.02

0.03

Rs (

m3 /C

)

Temperature (K)

(a)

Fig. 3.15. (a) Anomalous Hall co-efficient (Rs), evaluated from Fig. 3.13(a), plot with temperature for the Zn(Fe,Al)O thin films with 5% Fe doping. (b) Relation between Rs with linear resistivity of the film. The red line is the best fit curve with the Eq. 3.16 and the green line is the fitted curve with skew scattering . Red line fit describes well the Rs vs. ρ behavior of this Zn(Fe,Al)O DMS film with 5 % Fe doping.


73

Fig. 3.15(b) shows the Rs vs ρ plot. The fitted curve using 2ρρ sjsks baR += has been

shown by red line and 2ρsjs bR = has been shown by green line. The red curve fits very

well with the experimental data as shown in Fig. 3.15(b) and it confirms that both the

mechanisms are presents in the Zn(Fe,Al)O with 5% Fe film. The negative ask value (~-

268 units) is widely observed in ferromagnetic Transition Metals.

3.3.6.3. Magnetoresistance behaviors

Having a single-valley conduction band and the possibility of heavy n-type carrier

doping, ZnO is a suitable material for the study of magnetotransport. Isothermal MR has

been measured at different temperatures with the magnetic field parallel to the c-axis of

the films. Figure 3.16(a) and Fig. 3.16(b) are the nature of % of MR of the 5% Fe doped

Zn(Fe)O and Zn(Fe,Al)O films measured at different temperatures. Figure 3.16(c) and

Fig. 3.16(d) are the same for those films at lower temperature ranges. The magnetic field

-8 -6 -4 -2 0 2 4 6 8

-1

0

1

2

3

4 1.6 K 10 K 20 K 30 K 40 K 50 K 60 K 70 K 80 K 90 K 100 K

% M

R

Magnetic field (T)

(a)

Zn(Fe)O

-8 -6 -4 -2 0 2 4 6 8-6

-4

-2

0

2

4

6

2 K 10 K 20 K 30 K 40 K 50 K 60 K 70 K 80 K 90 K 100 K

% M

R

Magnetic field (T)

(b)

Zn(Fe,Al)O

-10 -8 -6 -4 -2 0 2 4 6 8 10-1.0-0.50.00.51.01.52.02.53.03.54.0

% M

R

Magnetic field (T)

1.6 K 2.3 K 4.2 K 6.2 K 8.4 K 9.4 K 10 K

Zn(Fe)O

(c)

-10 -8 -6 -4 -2 0 2 4 6 8 10-6

-4

-2

0

2

4

6

% M

R

Magnetic field (T)

1.6 K 2.3 K 4.2 K 6.2 K 8.4 K 9.4 K 10 K

Zn(Fe,Al)O (d)

Fig.3.16. (a) and (b) The magnetic field dependent %MR measured at different isothermal temperature of Zn(Fe)O and Zn(Fe,Al)O with 5% Fe concentration, respectively. (c) and (d) The %MR plot with magnetic field measured at lower temperatures upto 10 K for those same samples, respectively.


Chapter 3

74

dependent % of MR plot shows different behavior at different temperatures for all the

doped films as shown in Fig. 3.17.

The magnetoresistance (MR) of the films have been measured at different temperatures to

investigate the s–d exchange interaction between the conducting s-electron spins and the

d-electron spins localized at the magnetic Fe impurities [55]. We have observed the

positive MR at low magnetic field and negative MR at higher magnetic field. The s-d

exchange-induced spin splitting of the conduction band could account for positive MR

while suppression of electrons at weak localization of impurity centers could account for

the negative MR of the iron doped ZnO films. The behavior of MR at different field

range can be described in different four ways (i) A positive MR at lower field range

which arises in a two-band model from the action of the Lorentz force on the mobile

carriers. For carriers in closed orbits, this term is of the form 22

2%

cHbaHMR+

= . This

Fig. 3.17. %MR plot with magnetic field of Zn(Fe)O and Zn(Fe,Al)O thin films with different Fe concentrations. (a), (c) and (e) are the %MR nature of Zn(Fe)O thin films with different Fe concentrations at temperatures 1.6 K, 10 K and 20 K, respectively and (b), (d) and (f) are the %MR nature of Zn(Fe,Al)O samples at temperatures 1.6 K, 10 K and 20 K, respectively.


75

saturates MR at higher fields, (ii) the open orbits terms, for which the magnetoresistance

is quadratic and nonsaturating, is 2% dHMR = , (iii) A negative term arises due to spin-flip

scattering from singly occupied localized states with s =1/2 which is given by

1cosh%1

−⎟⎟⎠

⎞⎜⎜⎝

⎛=

−

fHMR , (iv) A negative term due to scattering from a paramagnetic or

ferromagnetic moment is given by 2% hMMR = where h is the dimensionless coefficient

which depends on the strength of the s-d scattering in the system. The parameters a, b, c,

d, f and h are the constants depend on the free carrier concentrations, magnetic properties

of the films etc.

The %MR is strongly temperature as well as free carrier concentration dependent.

The semiconducting transition metal doped ZnO films show low field positive %MR at

1.6 K as shown in Fig. 3.17(a) and (b). Increasing of carrier concentration by

incorporating 1% Al causes decrease of low field positive MR and increase of high field

negative MR in the films. MR is positive at 10 K as shown in Fig. 3.17(c) and (d). The

enhanced positive MR in the Al incorporated systems is caused by enhanced spin

splitting of conduction band due to s-d exchange interactions. At 20 K the MR is negative

at lower field and positive at higher field as shown in Fig. 3.17(e) and (f). MR shows

oscillatory behavior at that temperature. The negative MR decreases in the Al

incorporated films due to enhancement of spin splitting conduction band. The films show

Fig.3.18. The plot of %MR measured at 8 T magnetic field with temperatures of Zn(Fe)O and Zn(Fe,Al)O thin films with 5% Fe. Inset (a) and (b) are the same plot for 7% and 10% Fe doped films, respectively.

0 20 40 60 80 100

-6

-4

-2

0

2

4

6

Zn(FeO) Zn(Fe,Al)O

%M

R a

t 8 T

Temperature (K)

5%

0 20 40 60 80 100-2

-1

0

1

2

3

% M

R a

t 8 T

Temperature (K)

7%

0 20 40 60 80 100-0.9

0.0

0.9

%M

R a

t 8 T

Temperature (K)

10%

(a)

(b)


Chapter 3

76

negligible negative MR over 50 K as the insulating type DMS films shows. The

temperature dependent %MR has been shown in Fig. 3.18 for Zn(Fe)O and Zn(Fe,Al)O

films with 5% Fe. The insets (a) and (b) of Fig. 3.18 are the same plot for 7 and 10% Fe.

As temperature increases the resistivity due to spin disorder scattering increases and it

becomes larger than the resistivities due to impurity and thermal scattering. So, the MR

increases with increasing temperature. At higher temperature the MR starts decreasing

because of the both ionized impurity scattering and phonon scattering become

independent of sub-band spin spilt energy (δ), whereas spin-disordered scattering

decreases with δ. So, the MR vs. temperature curve shows the hump giving a peak of MR

at a certain temperature.

3.4. Summary

The high crystalline quality epitaxial Zn(Fe)O thin film doped with iron deposited

on sapphire substrate at a substrate temperature of 450 ºC shows room temperature

ferromagnetic behavior. Increasing of Fe doping concentration in the ZnO films

decreases the magnetic moment of the systems. The incorporation of 1% Al enhances the

saturation moment up to three times for 5% Fe doped samples. This result concludes that

there is the effect of free carrier density on magnetism. A clear correlation between the

magnetization per transition metal ion and the ratio of the number of carriers and number

of donors have been found in these films and establishes the theory of carrier induced

ferromagnetism. From the detailed low temperature magnetization investigation of all our

(Fe,Al) doped films it has been confirmed that the ferromagnetic moment along with a

huge paramagnetic components are present in each film. The steep exponential rise of

ferromagnetic moment and paramagnetic susceptibility of the films in the lower

temperature regimes do not follow the standard Curie law but it behaves more like to a

insulating type DMS. We have also attempted to establish the carrier induced effect in

our DMS films using defect model.

The magnetic and transport properties of the Zn(Fe)O and Zn(Fe,Al)O films with

5, 7 and 10% Fe concentration grown by a PLD technique are investigated. Temperature-

dependent Hall effect measurements have been performed on Zn(Fe)O and Zn(Fe,Al)O

highly crystalline epitaxial thin films. We have extracted the free carrier concentration


77

and Hall mobility using the ordinary Hall coefficient. The Hall data analysis revealed

that the dominant donor has an activation energy ranging from 33 to 41 meV. The

temperature dependent Hall mobility data has can be explained by several type of

scattering mechanism viz. ionized impurity, acoustic deformation, piezoelectric potential,

polar optical scattering and as well as grain boundary scattering. The high field saturating

nature of Hall resistivity in Zn(Fe,Al)O thin film also confirms the presence of

ferromagnetism in the film. The behavior of anomalous Hall co-efficient with linear film

resistivity confirms the presence of both scattering mechanisms (skew scattering and side

jump mechanism) in the Zn(Fe,Al)O epitaxial thin film. We have observed the positive

MR at low magnetic field and negative MR at higher magnetic field for all the doped

DMS fims. The s-d exchange-induced spin splitting of the conduction band could account

for positive MR while suppression of electron at weak localization of impurity centers

could account for the negative MR of the iron doped ZnO. Negative magnetoresistance at

higher magnetic field has been observed in transition-metal-doped ZnO DMS films at

lower temperature.

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[42] Qingyu Xu, Lars Hartmann, Heidemarie Schmidt, Holger Hochmuth, Michael Lorenz, Rüdiger Schmidt-Grund, Chris Sturm, Daniel Spemann, and Marius Grundmann, Metal-insulator transition in Co-doped ZnO: Magnetotransport properties, Phys. Rev. B 73, 205342 (2006). [43] S. S. P. Parkin, N. More, and K. P. Roche, Oscillations in exchange coupling and magnetoresistance in metallic superlattice structures: Co/Ru, Co/Cr, and Fe/Cr, Phys. Rev. Lett. 64, 2304 (1990). [44] Jagadeesh S. Moodera, Janusz Nowak, and Rene J. M. van de Veerdonk, Interface Magnetism and Spin Wave Scattering in Ferromagnet-Insulator-Ferromagnet Tunnel Junction, Phys. Rev. Lett. 80, 2941 (1998). [45] T. Ditel, J. Spalek, Effect of thermodynamic fluctuations of magnetization on the bound magnetic polaron in dilute magnetic semiconductors, Phys. Rev. Lett. 28, 1548 (1983). [46] Jing Wang, Zhengbin Gu, Minghui Lu, Di Wu, Changsheng Yuan, Shantao Zhang, Yanfeng Chen, Shining Zhu, and Yongyuan Zhu, Giant magnetoresistance in transition-metal-doped ZnO films, Appl. Phys. Lett. 88, 252110 (2006). [47] K T Roro, G H Kassier, J K Dangbegnon, S Sivaraya, J E Westraadt, J H Neethling, A W R Leitch and J R Botha, Temperature-dependent Hall effect studies of ZnO thin films grown by metalorganic chemical vapour deposition, Semicond. Sci. Technol. 23, 055021 (2008). [48] P. P. Debye and E. M. Conwell, Electrical Properties of N-Type Germanium, Phys. Rev. 93, 693(1954). [49] J. M. Dorkel, P. Leturcq, Carrier mobilities in silicon semi-empirically related to temperature, doping and injection level, Solid State Electronics 24, 821 (1981). [50] Seong-Il Kim, Chang-Sik Son, Min-Suk Lee, Yong Kim, Moo-Sung Kim, and Suk-Ki Min, Temperature dependent electrical properties of heavily carbon-doped GaAs grown by low-pressure metalorganic chemical vapor deposition, Solid State Comm. 93, 939 (1995). [51] Y. Z. Peng, T. Liew, T. C. Chong, C. W. An, and W. D. Song, Anomalous Hall effect and origin of magnetism in Zn1−xCoxO thin films at low Co content, Appl. Phys. Lett. 88, 192110 (2006). [52] Qingyu Xu, Lars Hartmann, Heidemarie Schmidt, Holger Hochmuth, Michael Lorenz, Rüdiger Schmidt-Grund, Daniel Spemann, and Marius Grundmann, Magnetoresistance effects in Zn0.90Co0.10O films, J. Appl. Phys. 100, 013904 (2006). [53] P. Khatua, T. K. Nath, and A. K. Majumdar, Extraordinary Hall effect in self-assembled epitaxial Ni nanocrystallites embedded in a TiN matrix, Phys. Rev. B 73, 064408 (2006). [54] Z. Yang, W. P. Beyermann, M. B. Katz, O. K. Ezekoye, Z. Zuo, Y. Pu, J. Shi, X. Q. Pan, and J. L. Liu, Microstructure and transport properties of ZnO:Mn diluted magnetic semiconductor thin films, J. Appl. Phys. 105, 053708 (2009). [55] P. Stamenov, M. Venkatesan, and L. S. Dorneles, D. Maude, and J. M. D. Coey, Magnetoresistance of Co-doped ZnO thin films, J. Appl. Phys. 99, 08M124 (2006).

Chapter 4

Junction magnetoresistance of Pt/Zn(Fe)O and

Pt/Zn(Fe,Al)O metal-dilute magnetic semiconductor

junction


International journal 1. Room temperature enhanced positive magnetoresistance in Pt and carrier induced Zn(Fe)O and Zn(Fe,Al)O

dilute magnetic semiconductor junction) by S. Chattopadhyay, T. K. Nath Journal of Applied Physics vol. 108, pp. 083904 (2010). Selected for Virtual Journal of Nanoscale Science & Technology for the October 25, (2010)

Conference/Symposia 2. Room temperature magnetic sensors with Zn(FeAl)O by Pt Schottky contact by S. Chattopadhyay, T. K. Nath

54th DAE Solid State Physics Symposium (2009)

Junction magnetoresistance of Pt/Zn(Fe)O and Pt/Zn(Fe,Al)O metal‐dilute magnetic semiconductorjunction

Chapter 4

81

4.1. Introduction

Spin-polarized electron injection into semiconductors has been a field of growing

interest in present microelectronics era. The injection and detection of a spin-polarized

current in a semiconducting material could combine magnetic storage of information with

electronic readout in a single semiconductor device [1]. Spin injection across a

ferromagnet-nonmagnetic metal interface provided a cornerstone for the field of spin

dependent transport in metals. The spin degree of freedom holds promise for the

realization of enhanced or novel device concepts and applications in microelectronics.

Most of the existing spintronic applications are based on metallic devices such as spin

valves [2], magnetic tunnel junctions [3], spin torque effects [4], domain wall devices [5],

etc. An important hurdle in this context is the inefficient injection of spin-polarized

currents from metallic ferromagnets into semiconductors due to the large mismatch in

conductivities [6]. Another research direction is the study of spin injection and transport

in more traditional devices aimed at room-temperature operation [7].

Zinc oxide based materials have many interesting and useful properties in the

field of optoelectronics and sensing devices [8,9]. Applications of such materials are

especially attractive on consideration of the low cost and lack of toxicity of zinc oxide.

The investigation on diluted magnetic semiconductors (DMS) [10-13] demonstrated an

application of ZnO as a host material for spintronic devices, which make use of electron

spin for data reading and writing. To integrate the DMS into present electronics, low-

dimensional structures are required for exploiting the advantages offered by the spin [14].

The basic DMS property of ZnO is that it shows ferromagnetism at room temperature.

In this chapter, the room temperature J-V properties of the junction between

paramagnetic novel metal, Pt and dilute magnetic semiconductor with 5, 7 and 10% iron

doped ZnO have been studied. In this work we have showed that the reasonably high

value of positive magnetoresistances persists at the junction at room temperature and it

depends on the magnitude of the magnetic moment of the dilute magnetic

semiconducting (DMS) ZnO films.

4.2. Experimental procedure

The detailed preparation method of the pulsed laser deposited iron doped epitaxial

ZnO thin films [Zn(Fe)O] and 1% Aluminum incorporated iron doped ZnO films

Chapter 4 Junction magnetoresistance in Pt and carrier induced Zn(Fe)O and Zn(Fe,Al)O junction

82

[Zn(Fe,Al)O] with iron concentrations 5, 7, 10% have been discussed in chapter-3

section-3.2. The point contact Pt metal has been used with the film to make a non-ohmic

(Schottky) type contact. Undoped ZnO thin film has also been grown on c-plane (0001)

sapphire substrate using the same PLD technique employing KrF excimer laser (λ = 248

nm) as discussed in chapter-3 for comparative study.

The I-V characterizations have been carried out using Keithley 2612 source meter

with 1 microvolt resolution. The magnetic field was applied in the direction of current

parallel to the film plane geometry using a high precision electromagnet (polytronic,

model HEM 100).

4.3. Results and discussion

4.3.1. Structural properties

The cross sectional high resolution transmission electron microscope (HRTEM)

image, shown in Fig. 4.1(a) clearly establishes that the ZnO films grown on (0001)

sapphire substrate are highly epitaxial and well crystalline in nature. The film - substrate

interface is very sharp having extremely good lattice matching between substrate and film

as discussed detailed in chapter-3.

(b)(a)

(c) (d)

Fig.4.1. (a) Cross sectional HRTEM image of Zn(Fe)O on sapphire substrate confirms that the films are epitaxial and well crystalline in nature. (b) Room temperature near edges EXAFS spectra for both Zn(Fe)O and Zn(Fe,Al)O epitaxial films compared to some standards. The valence looks to be Fe2+ for both the ZnO films. (c) and (d) are the AFM image of Zn(Fe)O for the scan area of 5 μm × 5 μm and 1 μm × 1 μm, respectively.


Chapter 4

83

The room temperature near edge EXAFS spectra for both the Fe and Fe with Al doped

epitaxial films along with some standards (FeO, Fe2O3 maghemite and Fe2O3 hematite)

have been shown in Fig. 4.1(b). The valence of Fe in both the DMS films appears to be

Fe2+ as the band edge positions are similar to FeO spectra. The small pre-edge peak of the

films is likely due to the less symmetric environment in the Zn site compared to the

octahedral coordination in FeO. The AFM image of the films recorded for 5 µm × 5 µm

and 1 µm × 1 µm scan area have been shown in Fig. 4.1(c) and Fig. 4.1(d), respectively.

The measured r.m.s. roughness of the films have been obtained to be ~ 1 nm. The

thicknesses of the films are about 0.3 to 0.4 μm. As discussed in earlier chapter all the

Zn(Fe)O and Zn(Fe,Al)O are epitaxial and highly crystalline in nature.


The room temperature ferromagnetic M(H) behavior of the Zn(Fe)O and

Zn(Fe,Al)O films grown on sapphire substrate at optimum deposition condition in the

magnetic field range of 0 to ± 1 T using a SQUID magnetometer have been shown in

Fig. 4.2.

-10000 -5000 0 5000 10000-0.4

-0.3

-0.2

-0.1

0.0

0.1

0.2

0.3

0.4

Zn(Fe)O

Zn(FeAl)O

H (Oe)

M (μ

B/F

e2+ )

(b)

-10000 -5000 0 5000 10000-0.15

-0.10

-0.05

0.00

0.05

0.10

0.15

5% Fe doped 7% Fe doped 10% Fe doped

M (μ

B/F

e2+ )

H (Oe)

(a)

-10000 -5000 0 5000 10000

-0.08

-0.04

0.00

0.04

0.08

Zn(Fe)O

Zn(FeAl)O

M (μ

B/F

e2+ )

H (Oe)

(c)

-10000 -5000 0 5000 10000-0.04

-0.03-0.02

-0.010.000.01

0.020.03

Zn(Fe)O

Zn(FeAl)O

H (Oe)

M (μ

B/F

e2+ )

(d)

Fig.4.2. Ferromagnetic M-H loop for all the Zn(Fe)O and Zn(Fe,Al)O films at room temperature. (a) M-H loop of Zn(Fe)O films with different Fe doping concentrations. (b), (c) and (d) are the comparative magnetization behavior between Zn(Fe)O and Zn(Fe,Al)O films with 5, 7 and 10% Fe concentrations, respectively. Al content is 1% for all the films.


84

The diamagnetic contributions of sapphire substrate have been subtracted carefully at

each magnetic field from the net magnetization (uncorrected raw data) to estimate the

actual ferromagnetic contribution of such ferromagnetic films at 300 K [Detailed work on

this is discussed in chapter-3]. Figure 4.2(a) shows the M-H loop of Zn(Fe)O samples

with different doping concentrations of iron. Figure 4.2(b), (c) and (d) show comparative

study of Zn(Fe)O and Zn(Fe,Al)O films with 5, 7 and 10% Fe doped, respectively. After

correcting the substrate contributions in SQUID raw data a ferromagnetic hysteretic

M(H) behavior at room temperature is observed for both kind films. The coercive field

and saturation magnetization for the Zn(Fe)O sample are found to be 135 Oe and 0.18

μB/Fe2+. In the case of 1% Al incorporated film the coercive field is 69 Oe and the

saturation magnetization is strikingly enhanced to 0.4 μB/Fe2+. The 7% and 10% Fe doped

Zn(Fe)O samples show magnetic moment 0.04 μB/Fe2+ and 0.02 μB/Fe2+ respectively.

The decrease of ferromagnetic moment with increasing concentration of iron may be due

to the increase of antiferromagnetic coupling between Fe pairs in the matrix. With

increase in the Fe doping in ZnO, the average distance between adjacent Fe2+ ions

reduces. As the antiferromagnetic energy is less than ferromagnetic energy, the

antiferromagnetic coupling between Fe2+⎯Fe2+ ions dominates at higher Fe

concentrations and act as a ferromagnetic moment killer reducing average magnetic

moment per Fe ion. Similar results are obtained for Mn doped and Ni doped ZnO films

[15,16]. 1% Al incorporation for those higher doping (7% and 10% Fe doping) cases also

enhances the magnetic moment mainly due to enhanced carrier induced ferromagnetism

[12].

4.3.3 Current-voltage characteristics without applied magnetic field

The current density-voltage (J-V) characteristics of Zn(Fe)O and Zn(Fe,Al)O with

Pt non-ohmic point contact with 0.28±0.01 mm2 contact area has been shown in Fig. 4.3.

Figure 4.3(a) shows the J-V behavior of Zn(Fe)O films with different Fe doping

percentages. Figure 4.3(b), (c) and (d) are the comparative J-V behavior of ZnO, Zn(Fe)O

and Zn(Fe,Al)O for 5%, 7% and 10% iron doping, respectively. The junction J-V

characteristics are denoted as [17],

⎟⎟⎠

⎞⎜⎜⎝

⎛ −=

kTIRVe

JJ s

η)(

exp0 (4.1)


Chapter 4

85

where, J0 is the reverse saturation current density. η and Rs are, ideality factor and

junction series resistance, respectively. The parameters evaluated from the forward J-V

curves of all the films shown in Fig. 4.3 have been summarized in Table 4.1. The ideality

factor of all Zn(Fe,Al)O samples lies between 1 and 2 and the values are near to 1 implies

that the thermionic emission dominates in Zn(Fe,Al)O samples whereas for Zn(Fe)O

films recombination degeneration transport process along with other defect induced

transport mechanism dominates.

4.3.4. Current-voltage characteristics with applied magnetic field

Figure 4.4 shows the forward J-V characteristics of such magnetic

semiconducting thin film junctions with Pt point contact and the J-V behaviors show

reasonably high sensitivity under magnetic field according to their magnetic moments.

Figure 4.4(a), (b) and (c) are the J-V plot of Zn(Fe)O films with 5, 7 and 10% iron

-8 -6 -4 -2 0 2 4 6 810-4

10-3

10-2

10-1

100

101

(b)Cur

rent

den

sity

(A/c

m2 )

ZnO Zn(FeAl)O Zn(Fe)O

Voltage (V)

5% Fe doped

-8 -6 -4 -2 0 2 4 6 810-4

10-3

10-2

10-1

100

101

(c)

Cur

rent

den

sity

(A/c

m2 )

ZnO Zn(Fe)O Zn(FeAl)O

Voltage (V)

7% Fe doped

-8 -6 -4 -2 0 2 4 6 810-4

10-3

10-2

10-1

100

101

(d)

Cur

rent

den

sity

(A/c

m2 )

Voltage (V)

ZnO Zn(Fe)O Zn(FeAl)O

10% Fe doped

-8 -6 -4 -2 0 2 4 6 810-4

10-3

10-2

10-1

100

Cur

rent

den

sity

(A/c

m2 )

Voltage (V)

ZnO Zn(Fe)O 5% Fe Zn(Fe)O 7% Fe Zn(Fe)O 10% Fe

(a)

Fig. 4.3. Junction J-V characteristics for all the Zn(Fe)O and Zn(Fe,Al)O films at room temperature. (a) J-V characteristics of Zn(Fe)O films with different Fe doping concentrations. (b), (c) and (d) are the comparative J-V study between Zn(Fe)O and Zn(Fe,Al)O films with 5, 7 and 10% Fe concentrations, respectively. Al content is 1% for all films.


86

doping, respectively, and Fig. 4.4(d), (e) and (f) are the same plots of Zn(Fe,Al)O

samples. The change of junction magneto-resistances (JMR) at a fixed bias voltage (7 V)

with applied magnetic field up to 0.6 T of different Zn(Fe)O and Zn(Fe,Al)O have been

shown in corresponding insets of Fig. 4.4. The J-V characteristics under magnetic field

have been fitted by the Eq. (4.1) and the parameters have been summarized in Table 4.1.

The series resistance increases with applied magnetic field and it shows the positive

junction magneto-resistance behavior of the films.

-0.6 -0.3 0.0 0.3 0.6

0

2

4

6

8

10

0 1 2 3 4 5 6 7 80

2

4

6

8

10

12

% J

MR

Magnetic field (T)

(e)

Cur

rent

den

sity

(A/c

m2 )

0.6 T

0 T

Voltage (V)

Zn(FeAl)O with 7% Fe and1% Al

-0.6 -0.3 0.0 0.3 0.6

0

2

4

6

8

0 1 2 3 4 5 6 7 80

2

4

6

8

10

12

% J

MR

Magnetic field (T) (f)

Cur

rent

den

sity

(A/c

m2 )

0.6 T

0 T

Voltage (V)

Zn(FeAl)O with 10% Fe and 1% Al

0 2 4 6 80

2

4

6

8

10

12

-0.6 -0.3 0.0 0.3 0.6

0

5

10

15

20

(d)C

urre

nt d

ensi

ty (A

/cm

2 )

0.6 T

Voltage (V)

0T

Zn(FeAl)O with 5% Fe and 1% Al

% J

MR

Magnetic field (T)

0 2 4 6 8 100.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

-0.6 -0.3 0.0 0.3 0.6

0

2

4

6

8

(b)

Cur

rent

den

sity

(A/c

m2 )

0.6 T

Voltage (V)

Zn(Fe)O with 7% Fe

0 T

Magnetic field (T)

% J

MR

-0.6 -0.3 0.0 0.3 0.60

1

2

3

0 2 4 6 8 100.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

% J

MR

Magnetic field (T) (c)

Cur

rent

den

sity

(A/c

m2 )

0.6 T0 T

Voltage (V)

Zn(Fe)O with 10% Fe

0 2 4 6 8 100.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

-0.6 -0.3 0.0 0.3 0.6

0

2

4

6

8

10

Zn(Fe)O with 5% Fe

Cur

rent

den

sity

(A/c

m2 )

Voltage (V)

0.6 T

0 T

(a)

% JM

R

Magnetic field(T)

Fig. 4.4. Junction J-V characteristics with and without applied magnetic field at room temperature; (a), (b) and (c) are the J-V properties of Zn(Fe)O film with Fe concentration 5, 7 and 10%, respectively. (d), (e) and (f) are the same of Zn(Fe,Al)O with Fe concentration 5, 7 and 10%, respectively. Insets of all the figures of Fig.4 are the plot of %JMR with different applied magnetic field of the corresponding junctions. The blue lines are the corresponding fitted curves.


Chapter 4

87

Table 4.1. Parameters extracted from the fitting of J-V characteristics

Applied

magnetic

field

Samples Reverse

saturation current

density (A/cm2)

Ideality factor Series

resistance

(Ω-cm)

0 T Zn(Fe)O with

5% Fe

0.025 6.89 2.67733

Zn(FeAl)O with

5% Fe

0.01429 1.12 0.61597

Zn(Fe)O with

7% Fe

0.04643 9.19 2.74142

Zn(FeAl)O with

7% Fe

0.01786 1.13 0.63501

Zn(Fe)O with

10% Fe

0.04643 9.19 2.82621

Zn(FeAl)O with

10% Fe

0.01429 1.13 0.66354

0.6 T Zn(Fe)O with

5% Fe

0.03929 10.72 2.81462

Zn(FeAl)O with

5% Fe

0.325 9.47 0.66514

Zn(Fe)O with

7% Fe

0.05714 12.54 2.7769

Zn(FeAl)O with

7% Fe

0.02143 2.51 0.68197

Zn(Fe)O with

10% Fe

0.04286 11.23 2.85482

Zn(FeAl)O with

10% Fe

0.03571 0.64 2.67733


88

4.3.5. Junction magneto-resistance properties

The increase of junction series resistance can be explained by the theoretical

model of spin tunneling in ferromagnetic to non-magnetic junctions [18]. The spin

injection process alters the potential drop across the F/N interface because differences of

spin dependent electrochemical potentials on either side of the interface generate an

effective resistance Rδ . It follows that 2/)0()0()0( sFFFnJ PJR μμμ σ+−= . R is the

junction series resistance, μn(0) and μF(0) are the electrochemical potentials for non-

magnetic and ferromagnetic sides of the junctions, J is the junction current density. PσF is

related to the conductivity polarization at the ferromagnetic interface and μsF(0) is the

spin accumulation at ferromagnetic side. Under magnetic field the junction series

resistance can be modified by RRR JJm δ+= where δR is the change of junction series

resistance and it can be expressed by,

( ) ( )

⎥⎥⎦

⎤

⎢⎢⎣

⎡ −++= ΣΣ

FN

FcFcFFN

rPPrrPrPrr

R222

σσδ (4.2)

where, rF, rN and rc are the ferromagnetic, non-ferromagnetic and contact resistance

respectively. ΣP is the contact conductivity polarization. rFN is the effective equilibrium

resistance of the Ferromagnetic/Non-ferromagnetic junction. From Eq. (4.2) it can be

clearly seen that the δR is always positive i.e. δR > 0 at higher applied potentials. The

positive junction MR and rectifying behavior has also been observed in ZnO

heterostructures with other ferromagnetic systems. Similar spin injection theory has been

evoked to explain their observed positive junction MR at the ferromagnet/semiconductor

interface [19,20].

The plot of % of junction MR (JMR) with applied magnetic field at a bias voltage

of 7 V has been shown in the insets of Fig. 4.4. It shows that the JMR behavior follows a

simple empirical relation with magnetic field as [21], βαHJMR= (4.3)

where, α and β are coefficients and are evaluated employing a non-linear least square

fitting using χ2 minimization technique. The coefficients thus obtained are listed in


Chapter 4

89

Table- 4.2 for all the samples. The coefficient β is lower than one at room temperature

showing nonlinear magnetic field dependence of positive MR of the junction.

Table 4.2. Fit parameters α and β from junction magnetoresistance plot

Sample α (T- β) β

Zn(Fe)O with 5% Fe 0.13 0.56

Zn(Fe)O with 7% Fe 10.31 0.70

Zn(Fe)O with 10% Fe 4.21 0.63

Zn(FeAl)O with 5% Fe 30.39 0.83



The % of JMR is found to be strongly dependent on the magnetic moments of the

respective magnetic semiconducting films. From the magnetization as a function of the

carrier density (nc) obtained from room temperature Hall voltage measurements of all the

Zn(Fe)O and Zn(Fe,Al)O films with 5% iron, the maximum saturation magnetization

(Ms) at 300 K is observed for the films with optimized carrier density of nc/ni ≈ 0.4 as

shown in Fig. 4.5(a). The film growth conditions were changed systematically (varing

oxygen pressure, laser energy, and target to substrate distance etc.) optimizing the best

DMS film property. The localized spins of the Fe ions are interacting with band electrons

and the standard theory of DMS can be applied. The % of JMR of those films is observed

to follow interestingly the same trend as the magnetic moment of the films follows

[shown in Fig. 4.5(b)]. With increasing doping percentage of iron the magnetic moment

decreases. Fig. 4.5(c) shows the plot of magnetic moment as a function of doping

concentration. The drop of moment with increasing iron concentration may be due to the

increasing of antiferromagnetic coupling between Fe pairs which occurs at shorter

separation distances. The % of JMR also decreases with increasing doping concentration

mimicking the same trend as the magnetic moments of the DMS films demonstrate.


90

4.4. Summary

The room temperature ferromagnetic iron doped ZnO with Pt metal point contact

shows non-ohmic J-V behavior at room temperature and shows reasonably high

sensitivity under magnetic field. The Pt/Zn(Fe)O junction shows positive junction

magnetoresistance at room temperature and the phenomenon can be best explained using

usual ferromagnetic to paramagnetic spin injection theory. Incorporation of 1% Al shows

higher junction magnetoresistance compared to without Al doped films. The junction

magnetoresistances are found to strictly depend on the magnitude of magnetic moments

of the DMS films. The magnetic moment depends on the carrier density and also the JMR

depends on magnetic moment. As magnetic moment decreases due to higher

concentration of Fe, JMR also mimics the same behavior.

References

[1] P. R. Hammar, B. R. Bennett, M. J. Yang, and M. Johnson, Observation of Spin Injection at a Ferromagnet-Semiconductor Interface, Phys. Rev. Lett. 83, 203 (1999). [2] I. Appelbaum, D. J. Monsma, K. J. Russell, V. Narayanamurti, and C. M. Marcus, Spin-valve photodiode, Appl. Phys. Lett. 83, 3737 (2003).

0.0 5.0x1020 1.0x1021 1.5x1021 2.0x1021 2.5x1021

8

10

12

14

16

18

20

% J

MR

Carrier concentration (cm-3)0 1x1021 2x1021

0.00

0.05

0.10

0.15

0.20

0.25

0.30

0.35

0.40

M (m

B/F

e2+ )

Carrier concentration (cm-3)

5 6 7 8 9 102468

101214161820

Zn(Fe)O

Zn(Fe,Al)O

% J

MR

Fe doping percentage5 6 7 8 9 10

0.000.050.100.150.200.250.300.350.40

Zn(Fe)O

Fe doping percentage

M (μ

B/F

e2+ )

Zn(Fe,Al)O

(a) (b)

(c) (d)

Fig. 4.5. (a) Room temperature carrier induced ferromagnetism in Zn(Fe)O and Zn(Fe,Al)O films with Fe concentration 5%. (b) % JMR of the corresponding films. (c) Plot of magnetic moment with different doping concentrations of iron. (d) Corresponding JMR of those films at room temperature.


Chapter 4

91

[3] A. Kalitsov, M. Chshiev, I. Theodonis, N. Kioussis, and W. H. Butler, Spin-transfer torque in magnetic tunnel junctions, Phys. Rev. B 79, 174416 (2009). [4] F. Junginger, M. Kläui, D. Backes, U. Rüdiger, T. Kasama, R. E. Dunin-Borkowski, L. J. Heyderman, C. A. F. Vaz, and J. A. C. Bland, Spin torque and heating effects in current-induced domain wall motion probed by transmission electron microscopy, Appl. Phys. Lett. 90, 132506 (2007). [5] D. A. Allwood, G. Xiong, C. C. Faulkner, D. Atkinson, D. Petit, and R. P. Cowburn, Magnetic Domain-Wall Logic, Science 309, 1688 (2005). [6] S. H. Chun, S. J. Potashnik, K. C. Ku, P. Schiffer, and N. Samarth, Spin-polarized tunneling in hybrid metal-semiconductor magnetic tunnel junctions, Phys. Rev. B 66, 100408 (2002). [7] W. Van Roy, P. Van Dorpe, R. Vanheertum, P. J. Vandormael, and G. Borgh, Spin Injection and Detection in Semiconductors—Electrical Issues and Device Aspect, IEEE Trans. Electron Devices 54, 933 (2007). [8] X. J. Zheng, B. Yang, T. Zhang, C. B. Jiang, S. X. Mao, Y. Q. Chen, and B. Yuan, Enhancement in ultraviolet optoelectronic performance of photoconductive semiconductor switch based on ZnO nanobelts film, Appl. Phys. Lett. 95, 221106 (2009) [9] S. W. Fan , A. K. Srivastava and V. P. Dravid, Nanopatterned polycrystalline ZnO for room temperature gas sensing, Sens. Act. B: Chem. 144, 159 (2010). [10] T. Dietl, H. Ohno, F. Matsukura, J. Cibert, and D. Ferrand , Zener Model Description of Ferromagnetism in Zinc-Blende Magnetic Semiconductors, Science 287, 1019 (2000). [11] X. J. Liu, X. Y. Zhu, C. Song, F. Zeng and F. Pan, Intrinsic and extrinsic origins of room temperature ferromagnetism in Ni-doped ZnO films, J. Phys. D: Appl. Phys. 42, 035004 (2009). [12] A. J. Behan, A. Mokhtari, H. J. Blythe, D. Score, X-H. Xu, J. R. Neal, A. M. Fox, and G. A. Gehring, Two Magnetic Regimes in Doped ZnO Corresponding to a Dilute Magnetic Semiconductor and a Dilute Magnetic Insulator, Phys. Rev. Lett. 100, 047206 (2008). [13] S. J. Pearton, C. R. Abernathy, M. E. Overberg, G. T. Thaler, D. P. Norton, N. Theodoropoulou, A. F. Hebard, Y. D. Park, F. Ren, J. Kim, and L. A. Boatner, Wide band gap ferromagnetic semiconductors and oxide, J. Appl. Phys. 93, 1 (2003) [14] C. Ronning, P. X. Gao, Y. Ding, Z. L. Wang, and D. Schwen, Manganese-doped ZnO nanobelts for spintronics, Appl. Phys. Lett. 84, 78 (2004). [15] D. L. Hou, R. B. Zhao, Y. Y. Wei, C. M. Zhen, C. F. Pan, G. D. Tang, Room temperature ferromagnetism in Ni-doped ZnO films, Curr. Appl. Phys. 10, 124 (2010) [16] W. B. Mi, H. L. Bai, Hui Liu, and C. Q. Sun, Microstructure, magnetic, and optical properties of sputtered Mn-doped ZnO films with high-temperature ferromagnetism, J. Appl. Phys. 101, 023904 (2007). [17] A. Singh, A Datta, S. K. Das, and V. A. Singh, Generalized RKKY interaction and spin-wave excitations Ferromagnetism in a dilute magnetic semiconductor, Phys. Rev. B 68, 235208 (2003). [18] S. J. May and B. W. Wessels, High-field magnetoresistance in p-InMnAs/n-InAs heterojunctions, Appl. Phys. Lett. 88, 072105 (2006).


92

[19] K. X. Jin, S. G. Zhao, C. L. Chen, J. Y. Wang, and B. C. Luo, Positive colossal magnetoresistance effect in ZnO/La0.7Sr0.3MnO3 heterostructure, Appl. Phys. Lett. 92, 112512 (2008). [20] S. Y. Park, Hyung Woo Lee, Young Soo Lee, D. F. Wang, Y. P. Lee and J. Y. Rhee, Magneto-transport properties of ZnO/La0.7Sr0.3MnO3 bilayer on p-Si(100), Phys. Stat. Sol. (c) 4, 4471 (2007). [21] T. Edahiro, N. Fujimura, and T. Ito, Formation of two-dimensional electron gas and the magnetotransport behavior of ZnMnO/ZnO heterostructure, J. Appl. Phys. 93,7673 (2003). [22] C. Song, X. J. Liu, F. Zeng, and F. Pan, Fully epitaxial ZnCoO/ZnO/ ZnCoO junction and its tunnel magnetoresistance, Appl. Phys. Lett. 91, 042106 (2007). [23] S. Honda, T. Ishikawa, K. Takai, Y. Mitarai, and H. Harad, New type magnetoresistance in Co/Si system, J. Magn. Magn. Mater. 290-291, 1063 (2005). [24] S. S. P. Parkin, C. Kaiser, A. Panchula, P. M. Rice, B. Hughes, M. Samant, and S. H. Yang, Giant tunnelling magnetoresistance at room temperature with MgO (100) tunnel barriers, Nature Mater. 3, 862 (2004). [25] J. Moser, M. Zenger, C. Gerl, D. Schuh, R. Meier, P. Chen, G. Bayreuther, W. Wegscheider, D. Weiss, C. H. Lai, R. T. Huang, M. Kosuth and H. Ebert, Bias dependent inversion of tunneling magnetoresistance in Fe/GaAs/Fe tunnel junctions, Appl. Phys. Lett. 89, 162106 (2006). [26] J. H. Hsua, S. Y Chen, W. M. Chang, C. R. Chang, Temperature dependence of magnetoresistance effect in Ag-Fe3O4 composites film, J. Magn. Magn. Mater. 272-276, 1772 (2004). [27] I. Žutić, J. Fabian and S. Das Sarma, Spintronics: Fundamentals and applications, Rev. Mod. Phys. 76, 323 (2004). [28] Z. G. Sheng, W. H. Song, Y. P. Sun, J. R. Sun, and B. G. Shen, Crossover from negative to positive magnetoresistance in La0.7Ce0.3MnO3-SrTiO3-Nb heterojunctions, Appl. Phys. Lett. 87, 032501 (2005).

Chapter 5

Structural, magnetic and electrical properties of

La0.7Sr0.3MnO3 thin films on p-Si


International journals 1. Low‐temperature resistivity minima in colossal magnetoresistive La0.7Sr0.3MnO3 thin film: A quantum

interference effect by S. Chattopadhyay and T. K. Nath, Solid state communications (Communicated)

Structural, magnetic and electrical properties of La0.7Sr0.3MnO3 thin films on p‐Si Chapter 5

93

5.1. Introduction

The perovskite manganites of the form R1-xAxMnO3 (R: rare earth elements, A: alkaline

earth elements) thin films have attracted much attention of the researchers due to its exceptional

electrical properties and the negative colossal magnetoresistance (CMR) effect [1-3]. The CMR

thin films have potential application in the field of the spintronics devices like magnetic field

sensor, hard disk read head and infrared bolometer. The large magnetoresistance (MR) ratio in

low magnetic field and at room temperature from the CMR materials have attracted much

attention in the area of research in manganites. There have been several reports on the La1-

xSrxMnO3 (LSMO) thin films [4-6] which attributes a high potential in application field.

The CMR effect and the correlated degrees of freedom of magnetic structure,

crystallographic structure and electrical resistivity in CMR materials, in addition to being of

fundamental scientific interest, appears to provide some scope for engineering in more sensitive

magnetoresistive response. The ‘colossal’ magnetoresistive (CMR) rare earth manganites display

a fascinating diversity of behaviors including several forms of magnetic, orbital and charge

ordering [7-9]. The materials also exhibit dramatic variations of physical properties with

frequency, temperature, chemical composition and applied strain, as well as the magnetoresistive

properties, which give them their colloquial name. The particular MR phenomena to be

described here are the massive decrease of resistance by application of a magnetic field [10-12].

The electronic inhomogeneities in the hole-doped La1-xSrxMnO3 have attracted considerable

attention with phase separation into conducting magnetic and insulating nonmagnetic domains

[13]. Among the La1-xSrxMnO3 families La0.7Sr0.3MnO3 shows the ferromagnetic behavior over

the room temperatures and a high value of magnetoresistance. Both the epitaxial and non-

epitaxial thin films of manganites show huge strain effects on their magnetic, electronic and

magneto-transport properties [14]. For device applications, it is necessary to find out the

properties of LSMO films on Si, as Si is widely used material in semiconductor industries. So, it

is necessary to find out the electrical and magnetic properties of the LSMO films with strain

effects. The defects in the crystal also affect the electrical and magneto-electrical properties and

hence it should be explored.

In this chapter, a detailed study of structural, magnetic, electrical, and magneto-electronic

properties of LSMO thin films have been explicitly studied. The effects of thickness and oxygen

Chapter 5 Structural, magnetic and electrical properties of La0.7Sr0.3MnO3 thin films on p‐Si

94

vacancy defects in the crystal have been elaborated and the dominating scattering process in the

conduction electron has been estimated.


5.2.1. Preparation of targets

The La0.7Sr0.3MnO3 powder have been synthesized through chemical pyrophoric reaction

process where we have employed stoichiometric mixtures of high purity La2O3 (99.99 %), SrCO3

(99.9+ %) and Mn(CH3COO)2 (99.0 %) [3]. After final grinding and pelletization of

La0.7Sr0.3MnO3 powders, the pelletized sample have been first heated at 800 °C for 12 h, then at

1000 °C for 12 h and at 1200 °C for another 12 h, with intermediate grinding. Final sintering of

the La0.7Sr0.3MnO3 target has been carried out at 1200 °C for 24 h.

5.2.2. Cleaning of substrates

La0.7Sr0.3MnO3 films on (100) p-Si substrate have been grown by Pulsed Laser Deposition

process using chemically synthesized single phase LSMO target. The substrates have been

cleaned in ultrasonic bath followed by chemical cleaning as mentioned below:

1. The p-Si (100) substrates have been first cleaned by de-ionized (DI) water in ultrasonic

chamber for 5 to 10 minutes to remove the dust particles.

2. Then the substrates have been cleaned ultrasonically for 5 to 10 minutes by acetone to

remove oils and greases over them.

3. Acetone has been cleaned using a high flow of DI water.

4. The mixture of NH4(OH), H2O2 and H2O with ratio 5:1:1 has been employed for substrate

cleaning to remove the acidic radicals and organic compounds. The substrates have been

kept in the solution till the end of reaction.

5. The substrates have been pulled out from the solution and cleaned using a high flow of DI

water.

6. Then the substrates have been kept in the mixture of H2O2 and H2SO4 (1:1) and boiled till

the boiling stops. It removes the basic radicals from the substrates and formed a SiO2 layer

over Si.

7. The substrates again cleaned using a high flow of DI water after removing it from the

solution.

8. Finally the oxide layers have been etched by dipping the substrate in 10 % HF solution.


95

Table 5.1: Condition for depositing different LSMO thin films on p-Si (100) substrates

Sample Name

Substrate treatment condition Film growth condition Sintering condition

T (ºC)

O2 pressure (mbar)

Oxidation time (min)

Final substrate Temp (ºC)

O2 pressure (mbar)

Deposition time (min)

Temp (ºC)

O2 pressure (mbar)

Sintering time (min)

Sample-1

800 10-5 No p-Si/SiO2 with native oxide 800 0.5 20 800 0.5 45

Sample-2

800 0.5 30 p-Si/SiO2 with thin oxide layer 800 0.5 20 800 0.5 45

Sample-3

800 0.5 45 p-Si/SiO2 with thick oxide layer 800 0.5 20 800 0.5 45

Sample-4

800 0.5 45 p-Si/SiO2 with thick oxide layer 800 5×10-3 20 800 5×10-3 45

Sample-5

800 0.5 45 p-Si/SiO2 with thick oxide layer 800 5×10-5 20 800 5×10-5 45

Sample-6


Sample-7



96

5.2.3. Deposition of La0.7Sr0.3MnO3 film

The films have been deposited on (100) p-Si substrates employing the pulsed laser deposition

technique using 248 nm KrF excimer pulsed mode laser. The different substrate treatment

conditions and film growth conditions for different LSMO thin films has been summarized in

Table 5.1. The substrate to target distance has been kept at 4 cm and the repetition rate of pulsed

laser (~ 10 pulses/second) has been used for all the films.

The electrical contacts have been made with high purity Ag-paste on LSMO film. The

temperature dependent electronic- and magneto-transport measurements have been carried out

using a source meter (Keithley, model - 2612), current source (Keithley, model-6221), PID

temperature controller (Lakeshore, model-331). A cryogen free ± 8 T superconducting magnet

with VTI system down to temperature 2 K (Cryogenics, U.K.) has been employed for high field

and low temperature transport measurements of these LSMO thin films.

5.3. Results and discussion

5.3.1. Structural study

The high resolution x-ray diffraction pattern (HRXRD) of LSMO film deposited on (100)

p-Si substrate using Cu-Kα radiation has been shown in Fig. 5.1. The multipeaks of LSMO

sample reveal the non-epitaxial nature of the LSMO film on SiO2/Si layer. Figure 5.1 (a) shows

the HRXRD pattern of LSMO film on Si/SiO2 substrate for different thickness (different pulse

duration at 0.5 mbar O2 and 800 ºC). Figure 5.1 (b) shows the HRXRD patterns of LSMO thin

films deposited in different O2 pressure.

Fig. 5.1. XRD pattern of LSMO thin film on Si/SiO2 up to 60º 2θ scan (a) for different pulse duration, (b) different O2 atmosphere

20 30 40 50 60

(211)(210)

(200)(111)

(110)(100) Sample-5

(b)

Sample-3

2θ (degree)

Sample-4

Inte

nsity

(a.u

.)

20 30 40 50 60

Sample-3

2θ (degree)

(210)

Sample-6

Inte

nsity

(a.u

.) (211)

(200)

(111)

(110)(100) Sample-7

(a)


97

All the films show crystalline growth in (100), (110), (111), (210) and (211) planes and

have little shift in 2θ towards higher angle than a bulk one implying that the strain affects the

crystal structure of non-epitaxial films. The shift in 2θ and full width half maximum (FWHM)

for different films have been summarized in Table 5.2.

Table 5.2: Peak position in 2θ and FWHM for different LSMO thin films

Sam

ple

FWHM 2θ

100 110 111 200 210 211 100 110 111 200 210 211

Sam

ple-3

0.15

47

0.21

043

0.30

978

0.27

78

0.25

898

0.43

56

23.02

566

32.78

6

40.49

781

46.96

067

52.95

644

58.50

877

Sam

ple-6

0.13

493

0.19

281

0.30

41

0.25

256

0.24

628

0.45

302

23.02

546

32.78

883

40.50

427

46.95

527

52.94

312

58.50

161

Sam

ple-7

0.17

711

0.25

515

0.29

465

0.31

115

0.38

229

0.39

15

23.12

937

32.92

348

40.62

412

47.08

569

53.09

857

58.67

025

Sam

ple-4

0.20

44

0.31

499

0.46

937

0.37

259

0.23

701

0.45

42

23.26

389

33.04

751

40.72

835

47.32

32

53.25

304

58.84

88

Sam

ple-5

0.20

26

0.26

553

0.31

6

0.30

433

0.20

203

0.44

257

22.89

26

32.67

681

40.33

589

46.83

358

52.72

188

58.38

022

The thin films grown on lower O2 atmosphere shows greater FWHM. The oxygen vacancy in the

films grown in O2 atmosphere causes a crystal deformation and can be a cause of higher strain in

the crystals. Due to increase of strain the FWHM also increases.

The preferred crystalline orientation of the nickel films has been evaluated by the texture

coefficients (TC) given by [15],

( )∑=

0

0

/1/

hklhkl

hklhkl

IIn

IITC (5.1)

Where, Ihkl and Iohkl are the diffraction intensity of the crystal plane (hkl) of the deposited and

bulk standard samples, respectively. n is the number of diffraction peak appeared in the HRXRD


98

pattern. If the texture coefficient is greater than 1.0, it indicates the existence of a preferred

orientation. The TCs found for different planes in different samples are summarized in Table 5.3.

Table 5.3. Texture coefficients of different planes in LSMO thin films

Sample TC

(100) (110) (111) (200) (210) (211)

Sample-7 3.32868 0.31607 0.23579 1.57143 0.34877 0.19925

Sample-6 3.45816 0.26569 0.17807 1.65635 0.24201 0.19972

Sample-3 3.91091 0.18001 0.12874 1.46965 0.17831 0.13239

Sample-4 3.29052 0.33756 0.12235 2.08297 0.08709 0.0795

Sample-5 3.4024 0.466 0.57652 0.75415 0.32213 0.4788

The texture coefficient is greater than 1.0 for the plane (100) for all the films implies that

the preferred orientation of LSMO thin films on (100) p-Si is crystallographic (100) plane. The

texture coefficient of (100) plane is higher for the film with higher thickness. The deficiency in

oxygen pressure causes a lower texture co-efficient of the films towards (100) plane.

(a) (b)

(c) (d)

Fig.5.2. FESEM image of the surface of LSMO films (a) Sample-3, (b) Sample-6, (c) Sample-4 and (d) Sample-5


99

5.3.2. Surface morphology

The surface morphology obtained from field emission scanning electron microscopy

images of different LSMO films have been shown in Fig. 5.2. Figure 5.2 (a), (b), (c) and (d) are

the FESEM micrograph of Sample-3, Sample-6, Sample-4 and Sample-5, respectively. The

images show the granular growth of the grains of the films. Figure 5.3 shows the cross sectional

FESEM micrograph of LSMO films deposited in 0.5 mbar O2 pressure and 800 °C substrate

temperatures for different pulse duration (growth time).

Figure 5.3 (a), (b) and (c) are the cross sectional FESEM image of Sample-3, Sample-6 and

Sample-7, respectively. All the films show rod like structure with nano dimensions. The plot of

thickness of LSMO films with pulse duration (growth time) have been shown in Fig. 5.3 (d). The

plot shows almost linear nature of thickness with pulse duration in this growth time range

confirming uniform growth of the films.

(a) (b)

(c)

642 nm 538 nm

448 nm

10 12 14 16 18 20

450

500

550

600

650

Thi

ckne

ss (n

m)

Pulesed duration (min)

(d)

Fig. 5.3. Cross sectional FESEM micrograph of LSMO films; (a) Sample-3, (b) Sample-6 and (c) Sample-7. (d) Plot of thickness with pulse duration (growth time).

(b) (a)

(c)


100


The magnetic properties of the films have been characterized using Quantum Design

ever-cool SQUID - VSM magnetometer. The temperature dependent magnetic M-H hysteresis

loops have been shown in Fig. 5.4 after correcting the substrate contribution as discussed in

chapter-3. The saturation fields are found to be ~ 0.07 T which is little greater than the bulk

sample [16]. This may due to the strain effect of non-epitaxial LSMO films. The clear hysteresis

at all temperatures up to 300 K shows the ferromagnetic behavior.

The temperature dependent field-cool (FC) and zero field cool (ZFC) magnetization

measurements have been carried out for those LSMO films to investigate the enhanced grain

surface effect of non-epitaxial thin films. The FC and ZFC properties at a constant magnetic field

0 50 100 150 200 250 3000.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8FC

Temperature (T)

mom

ent (

emu) 0.1 T

Sample-3

ZFC

Fig. 5.5. Temperature dependence of FC and ZFC magnetic measurement at 100 Oe for LSMO films

Fig. 5.4. Corrected temperature dependent M-H loop of the LSMO thin films on p-Si substrate (a) for sample-3 and (b) for sample-5

-4000 -2000 0 2000 4000

-60

-40

-20

0

20

40

60

(b)

Sample-5

5 K100 K200 K 300 K

Mag

netic

mom

ent (

emu/

gm)

Magnetic Field (Oe)-4000 -2000 0 2000 4000

-100

-50

0

50

100

Sample-3

5 K 50 K 100 K 300 K

Mag

netic

mom

ent (

emu/

gm)

Magnetic Field (Oe)

(a)


101

of 100 Oe has been shown in Fig. 5.5. The curves show strong irreversibility between FC and

ZFC curves exhibiting prominent maxima at Tmax (~ 130 K). The strong decrease of low field

ZFC and the increase of irreversibility with decreasing temperature can be attributed to the

passage from a ferromagnetic state to a low temperature disordered regime surface [17]. With

reducing temperature local anisotropy at the grain surface increases more sharply than the

exchange stiffness parameter due to lack of symmetry at grain surface. Consequently, opposing

magnetic interaction stabilized a spin glass like state at the grain surface, yielding freezing of

surface spin in random direction. This inhibits the exchange interaction to transmit across the

interface. But, they are expected to interact through dipolar interaction. The strength of

interaction depends on the assembly of grains which either can show superparamagnetic

blocking or cooperative freezing of moment at Tmax.

5.3.4. Electrical transport properties

The electrical characterization of all LSMO films having different thickness and different

oxygen deficiency has been carried out using a source meter (Keithley, model - 2612), current

source (Keithley, model -6221), PID temperature controller (Lakeshore, model-331). A cryogen

free 8 T superconducting magnet with VTI system using closed cycle helium refrigeration

technique down to temperature 2 K (Cryogenics, U.K.) has been employed for high field and low

temperature transport measurements.

5.3.4.1. ρ-T behavior without applied magnetic field

Figure 5.6 shows the temperature dependent resistivity measurements for different

LSMO thin films. Figure 5.6 (a) shows the ρ-T plot under 0 T applied magnetic field for LSMO

films with different thickness grown on SiO2/p-Si at a substrate temperature of 800 oC in 0.5

mbar O2 atmosphere. The absolute value of resistivity increases little with decreasing of

thickness of the films but the overall ρ-T behavior looks almost identical for all films. The films

show resistivity minima (TM) at low temperatures (<50 K). The double exchange interaction

model of Zener associated with the Jahn-Teller splitting of Mn d-levels and a strong electron-

phonon interaction explain most of the electrical and magnetic properties of these manganites

[18]. All the films show a metal-insulator transition temperature (Tp) at higher temperature.


102

According to the established phase diagram, a phase transition occurs from metal to insulator

phase giving a peak temperature (TP). The TM and TP does not change much with the variation of

thickness of the LSMO films.

Figure 5.6 (b) shows the ρ-T plot under 0 T applied magnetic field for LSMO films grown on

different oxygen ambient. The ρ-T behavior shows distorted nature with deficiency of oxygen.

The low temperature resistivity is much higher which may be due to the higher impurity

scattering in high oxygen deficient films. The curvature near TM and TP are much broader than

the well crystalline films. This may be due to creation of trap charges in disordered high oxygen

vacancy films.

5.3.4.2. ρ-T behavior under applied magnetic field

The temperature dependent resistivity plots with applied magnetic field up to 8 T for the

LSMO films with different thicknesses have been shown in Fig. 5.7. Figure 5.7 (a), (b) and (c)

are the ρ-T plots with different magnetic field for sample-3, sample-6 and sample-7, respectively.

Figure 5.7 (d), (e) and (f) are the magnetic field dependent magnetoresistance (MR) plots for the

corresponding films. Resistivity decreases drastically under applied magnetic field mainly due to

the spin polarized tunneling of mobile eg electrons across the ferromagnetic manganite grain

boundaries. The peak temperature shifts towards higher temperature with applied magnetic field.

The films shows negative magnetoresistance under applied magnetic field and it does not

Fig. 5.6. Temperature dependent resistivity plot of LSMO thin film (a) with different thickness, (b) with different oxygen vacancies

0 50 100 150 200 250 300

0.8

1.0

1.2

1.4

1.6 Sample-7

Sample-6

Res

istiv

ity (k

Ω-c

m)

Temperature (T)

Sample-3H=0

(a)

0 50 100 150 200 250 300

0.7

0.8

0.9

1.0

1.1

1.2

1.3

Sample-5

Res

istiv

ity (k

Ω-c

m)

Temperature (K)

Sample-3

Sample-4

(b)


103

changes much with the film thickness. Hwang et al. [19] has described a model for spin polarized

tunneling MR across the ferromagnetic manganite grain boundaries as,

3

0

)( KHJHdkkfAMRH

−−−= ∫ (5.2)

In zero field the domain boundaries are pinned at the grain boundary pinning centre with pinning

strength k. The grain boundaries have a distribution of pinning strength f(k) defined as the

minimum field needed to overcome a particular pinning barrier [20],

)exp()exp()( 222 DkCkBkAkf −+−= (5.3)

A, B, C, D, J and K are the parameters. The best fit values of the parameters have been listed in

Table 5.4 for all the LSMO films. The spin polarized tunneling MR ( ∫−=H

spt dkkfMR0

)( ) and

intrinsic MR (MRint) contributions have be evaluated using all the fit parameters keeping H = 8 T

for each LSMO films as listed in Table 5.4.

Fig. 5.7. Temperature dependent resistivity under applied magnetic field for (a) sample-3, (b) sample-6 and (c) sample-7; (d), (e) and (f) are corresponding % MR plot of the films with applied magnetic field.

0 50 100 150 200 250 3000.4

0.6

0.8

1.0

1.2

1.4

1.6

Res

istiv

ity (k

Ω-c

m)

Temperature (K)

Sample-6

0 50 100 150 200 250 3000.40.50.60.70.80.91.01.11.21.3

8 T7 T6 T5 T4 T3 T2 T1 T

Res

ista

nce

(kΩ

−cm

)

T(K)

0 T

Sample-3

0 2 4 6 8-50

-40

-30

-20

-10

0

2.8 K 5 K 10 K 50 K 100 K 150 K 200 K 250 K 300 K

% M

R

Magnetic Field (T)

0 50 100 150 200 250 3000.4

0.6

0.8

1.0

1.2

1.4

1.6

Res

istiv

ity (k

Ω-c

m)

Temperature (K)

Sample-7

0 2 4 6 8-50

-40

-30

-20

-10

0

2.8 K 5 K 10 K 50 K 100 K 150 K 200 K 250 K 300 K

% M

R

Magnetic field (T)0 2 4 6 8

-50

-40

-30

-20

-10

0 2.8 K 5 K 10 K 50 K 100 K 150 K 200 K 250 K 300 K

% M

R

Magnetic field (T)


104

Table 5.4. Evaluated fitting parameter using Eq. (5.2) and Eq. (5.3)

Sample T (K) A B C D J K MRspt(%) MRint(%)

Sample-3 2.8 -32.02 0.10 21.16 0.34 -4.21 -0.042 -42.3 -2.1

5 -33.15 0.15 21.86 0.39 -0.50 -0.011 -36.1 -9.1

10 -35.81 0.19 29.67 0.47 -1.32 -0.006 -32.0 -12.9

50 -30.92 0.19 24.63 0.47 -1.45 -0.006 -28.9 -13.8

100 -23.80 0.21 19.56 0.49 -2.39 -0.002 -20.7 -19.1

150 -18.99 0.21 14.80 0.49 -2.37 -0.002 -17.6 -20.1

200 -13.47 0.21 10.97 0.48 -3.07 0.0027 -11.4 -23.3

250 -7.93 0.21 5.80 0.48 -2.93 0.0012 -7.6 -22.8

300 -2.27 0.21 0.49 0.49 -5.20 0.0262 -3.7 -19.7

Sample-6 2.8 -38.13 0.22 24.5 0.48 -1.43 -0.004 -39.3 -4.1

5 -32.09 0.21 23.28 0.49 -1.34 -0.005 -31.9 -12.6

10 -34.03 0.21 28.21 0.49 -1.77 -0.001 -29.3 -14.4

50 -29.15 0.21 22.67 0.48 -1.70 -0.003 -26.1 -15.6

100 -21.95 0.22 16.66 0.49 -2.43 0.0003 -19.9 -18.7

150 -16.95 0.22 11.86 0.49 -2.47 -0.0005 -16.7 -19.8

200 -11.55 0.23 8.33 0.46 -3.08 0.004 -9.5 -24.1

250 -5.99 0.22 3.04 0.5 -2.85 0.0018 -7.5 -21.6

300 -2.63 0.22 0.47 0.49 -2.85 0.0031 -4.3 -17.8

Sample-7 2.8 -39.70 0.23 25.41 0.49 -2.18 0.0005 -40.5 -5.4

5 -32.53 0.22 23.63 0.49 -1.96 -0.0014 -30.9 -16.1

10 -33.15 0.22 26.05 0.50 -2.17 3E-5 -29.9 -16.2

50 -29.71 0.23 23.36 0.49 -2.37 0.001 -24.7 -19.5

100 -22.7 0.23 18.25 0.5 -3.06 0.004 -19.1 -22.3

150 -18.19 0.22 13.75 0.49 -2.87 0.001 -16.6 -22.7

200 -12.65 0.22 9.71 0.49 -3.44 0.005 -11.3 -25.1

250 -7.06 0.22 4.46 0.50 -3.32 0.004 -7.7 -24.4

300 -1.52 0.21 -0.34 0.49 -3.4 0.006 -2.5 -22.9


105

The % of MRspt drops sharply, where as the % of MRint increases with increasing temperature.

The temperature dependent % of MRspt can be explained using an empirical expression

( )/()( TcbaMRspt ++= [19].

5.3.4.2.1. ρ-T behavior below TM

Focusing on the resistivity plot again, it is necessary to study the appearance of low

temperature minima and peak temperature and their dependence on magnetic field. To explore

the fact we have chosen the ρ-T behavior of the sample-3. The low temperature ρ-T behavior

below resistivity minimum (TM) showing resistivity upturn in all LSMO thin films can have

various source of origin. At low temperatures, low resistive dilute alloys with very small

magnetic impurity generally show Kondo resistivity minima and the ρ(T) relation is given by

[21]

)ln()( 0 TCT −= ρρ (5.4) where ρ0 is the residual resistivity. The Kondo resistivity minima are attributed to the localized

magnetic impurities that are far apart and interact by polarizing electrons. The Kondo minima

disappear with the application of external magnetic field.

Moreover, in the disordered highly resistive systems the mean free path of conduction

electrons become small and they involve in multiple elastic scattering [22]. It causes a higher

resistivity in the system. Any inelastic process or electron-electron interaction or applied

magnetic field can reduce the resistivity. A resistivity minimum occurs at TM because there is

ultimately an increase of resistivity with temperature due to inelastic high temperature electron-

phonon scattering. Below the resistivity minima due to the electron-electron interaction the ρ(T)

relation is given by [23], 2/1

0)( BTT −= ρρ (5.5)

On the other hand, the resistivity due to the temperature dependent other scatterings like

electron-phonon, electron-magnon can be expressed as, nATT =)(ρ (5.6)

Figure 5.8 shows the low temperature minima of sample-3 with different applied

magnetic fields up to 8 T. The resistivity curves lower than TM (T < TM) have been fitted

considering both Kondo effect and electron-electron correlation effect as described in Eq. (5.4)


106

and Eq. (5.5). The temperature dependent inelastic term (Eq. 5.6) has been added with these

equations and hence we can write, nATTCT +−= )ln()( 0ρρ (5.7)

nATBTT +−= 2/10)( ρρ (5.8)

All the ρ(T) curves measured at different magnetic fields below TM have been fitted with

both Eq. (5.7) and Eq. (5.8) as shown in Fig. 5.8. Figure 5.8(a) is the low temperature ρ(T)

curves fitted with Eq. (5.8) (red lines) and Fig. 5.8(b) is the same fitted with Eq. (5.7) (green

lines). The best fit χ2 values for both the fits have been examined and it is found that the Eq. (5.8)

fits much better as compared to Eq. (5.7). Comparative % of deviation of fit with Kondo effect

and e-e interaction effect for the ρ(T) curves with 0 T and 8 T applied magnetic field have been

shown in Fig 5.8(c) and (d), respectively. It clearly shows that the e-e interaction model fits

much better thereby explaining best possible electronic transport mechanism at very low

temperature in these LSMO thin films. The evaluated parameters fitting with Eq. (5.8) have been

summarized in Table 5.5. From the Table 5.5 it is clear that the coefficient B (e-e interaction

term) in Eq. (5.8) remains unaltered with the application of magnetic field. However, the

coefficient A changes noticeably with the magnetic field keeping the value of the exponent n ~ 2

( ~ AT2). So the origin of this additional term in Eq.(5.8) possibly arises due to the electron-

magnon scattering process these films.

0 10 20 30 40 500.4

0.5

0.6

0.7

0.8

Res

istiv

ity (k

Ω-c

m)

Temperature (K)

Kondo minima

(b)

10 20 30 40 50

-0.8

-0.4

0.0

0.4

0.8

% o

f D

evia

tion

Temperature (K)

Kondo Effect

e-e interactionH=0

(c)

10 20 30 40 50

-0.2

-0.1

0.0

0.1

0.2

% o

f dev

iatio

n

Temperature (K)

Kondo Effect

e-e interaction

(d)

H=8 T

0 10 20 30 40 500.4

0.5

0.6

0.7

0.8

0.9

0T0T0T0T0T0T0T

Res

istiv

ity (k

Ω−c

m)

Temperature (K)

0T

e-e interation

(a)

Fig. 5.8. Low temperature ρ(T) curves fitted with (a) electron-electron interaction model, (b) Kondo Effect model; (c) and (d) are % of deviation of fit values at 0 and 8 T. The deviation clearly shows the good fit with e-e interaction model.


107

Table 5.5: Fit parameters evaluated using Eq. (5.8).

Magnetic field

(T)

ρ0 (kΩ-cm) B (kΩ-cmK-1/2) A ((kΩ-cmK-n)

×10-3

n

0 0.912 0.0385 0.742 1.82

1 0.675 0.0314 2.143 1.57

2 0.639 0.0282 1.343 1.66

3 0.613 0.0277 1.612 1.61

4 0.589 0.0272 2.061 1.55

5 0.567 0.0268 2.393 1.51

6 0.547 0.0259 2.396 1.50

7 0.529 0.0266 3.672 1.41

8 0.510 0.0256 3.589 1.41

The depth of minima )6.1(/)]()6.1([ KTK M ρρρ − has been plotted with magnetic field as

shown in Fig. 5.9. The depth of minima initially decreases slightly with the magnetic field till

2 T beyond which it becomes almost field independent.

5.3.4.2.2. ρ-T behavior above TM

The temperature dependent resistivity above TM and below TP (metal-insulator transition

temperature) as shown in Fig. 5.7(a) has been best described with the help of Matthiessen’s law

as described below [24]:

0 2 4 6 80.115

0.120

0.125

0.130

0.135

0.140

Dep

th o

f min

ima

Magnetic field (T)

Fig. 5.9. Magnetic field dependent


108

mlT ρρρρ ++= 0)( (5.9)

where, ρ0 is residual resistivity. The two other resistivity terms (ρl and ρm) originate from the

scattering of lattice phonon and spin wave at finite temperatures, respectively. For both lattice

and magnetic scatterings conduction electrons might undergo s-s and s-d transitions. The

scattering due to lattice have been given by [25],

∫ −−−⎟⎟⎠

⎞⎜⎜⎝

⎛=

T

xx

nn

Dl

D

eedxxTA

θ

θρ

0 )1)(1( (5.10)

where, the θD is the Debye temperature; n is a constant generally it is 3 for magnetic metal and

alloys with large d-band density of state.

The resistivity due to exchange interaction between the conduction electrons (s) and the

localized magnetic electrons (3d) is ρm. This interaction is generally called the s-d interaction.

This spin disorder resistivity goes with the 2nd power of temperature ( 2)( DTT =ρ ) where, D

shows the strength of s-d interaction. Eq.(5.9) fits well with the all ρ(T) curves in this

temperature regimes and describes well the electron transport mechanism in these LSMO films.

5.4. Summary

The LSMO thin films with different thicknesses and different oxygen vacancies have

been grown on SiO2/p-Si substrates where the SiO2 layer have been created by oxidation of Si.

All the films show nano rod like growth with high texturing towards (100) direction. The growth

of the film is almost linear with deposition time implying good uniformity of the films.

The magnetization measurement of the film shows good ferromagnetic nature at all

temperatures with saturation field near 0.07 T. The FC and ZFC curves show strong

irreversibility exhibiting prominent maximum at Tmax (~ 130 K). The sharp decrease of low field

ZFC curve at low temperature and the increase of irreversibility with decreasing temperature

generally are attributed to the passage from a ferromagnetic state to a low temperature disordered

surface regime.


109

The temperature dependent resistivity shows the metal-insulator transition giving a peak

temperature near 250 K. The peak temperature does not change, though the resistivity increases

slightly with decreasing thickness. The distortion in ρ-T curves for oxygen deficient films may be

due to the oxygen vacancy and lattice defects. The low temperature resistivity is much higher

which may be due to the higher impurity scattering in high oxygen deficient films. The curvature

near TM and TP are much broader than the well crystalline film possibly due to the creation of

trap charges in disordered high oxygen vacancy films. All the films show negative MR behavior

for temperatures up to 300 K and there is not much effect of film thickness in that thickness

range. Resistivity decreases sharply with the application of magnetic field in the low magnetic

field region mainly due to the spin polarized tunneling of mobile eg electrons across the

ferromagnetic manganite grain boundaries. The MR contribution due to extrinsic spin polarized

tunneling across the grain boundaries as well intrinsic CMR inside grains has also been separated

out for the films.

The resistivity behavior in the temperature range lower than TM has been well described

through the electron-electron interaction model rather than the Kondo effect. The dependency of

parameter A with applied magnetic field and the depth of minima indicate the electron-magnon

scattering process in the films. The resistivity behavior in the higher temperature regime (> TM)

has been described through the lattice and magnetic scattering using the Matthiessen’s law.

References

[1] P. Dey, T. K. Nath and A. Banerjee, Enhanced grain surface effect on magnetic properties of La0.5Gd0.2Sr0.3MnO3 nanoparticles: A comparison with bulk counterpart, Appl. Phys. Lett. 91, 012504 (2007). [2] P. Dey and T. K. Nath, Tunable room temperature low-field spin polarized tunneling magnetoresistance of La0.7Sr0.3MnO3 nanoparticles, Appl. Phys. Lett. 89, 163102 (2006). [3] P. Dey and T. K. Nath, Effect of grain size modulation on the magneto- and electronic-transport properties of La0.7Ca0.3MnO3 nanoparticles: The role of spin-polarized tunneling at the enhanced grain surface, Phys Rev. B 73, 214425 (2006). [4] J. Fontcuberta, M. Bibes, and B. Martı´nez, V. Trtik, C. Ferrater, F. Sa´nchez, and M. Varela, Tunable epitaxial growth of magnetoresistive La2/3Sr1/3MnO3 thin films, J. Appl. Phys. 85, 4800 (1999). [5] V. Moshnyaga, I. Khoroshun, A. Sidorenko, P. Petrenko, A. Weidinger, M. Zeitler, B. Rauschenbach, R. Tidecks, and K. Samwer, Preparation of rare-earth manganite-oxide thin films by metalorganic aerosol deposition technique, Appl. Phys. Lett. 74, 2842 (1999).


110

[6] A. Goyal, M. Rajeswari, R. Shreekala, S. E. Lofland, S. M. Bhagat, T. Boettcher, C. Kwon, R. Ramesh, and T. Venkatesan, Material characteristics of perovskite manganese oxide thin films for bolometric applications, Appl. Phys. Lett. 71, 2535 (1997). [7] A. Urushibara, Y. Moritomo, T. Arima, A. Asamitsu, G. Kido and Y. Tokura, Insulator-metal transition and giant magnetoresistance in La1-xSrxMnO3, Phys. Rev. B 51, 14103 (1995). [8] P. Schiffer, A. P. Ramirez, W. Bao, and S. W. Cheong, Low Temperature Magnetoresistance and the Magnetic Phase Diagram of La1-xCaxMnO3, Phys. Rev. Lett. 75, 3336 (1995). [9] Y. Tokura, Y. Tomioka, H. Kuwahara, A. Asamitsu, Y. Moritomo, and M. Kasai, Origins of colossal magnetoresistance in perovskite-type manganese Oxides, J. Appl. Phys. 79, 5288 (1996). [10] S. Jin, T. H. Tiefel, M. McCormack, R. A. Fastnacht, R. Ramesh and L. H. Chen, Thousand fold Change in Resistivity in Magnetoresistive La-Ca-Mn-O Films, Science 264, 413 (1994). [11] J. D. Boeck, Switching with Hot Spins, Science 281, 357 (1998). [12] G. A. Prinz, Magnetoelectronics, Science 282, 1660 (1998). [13] U. R. Singh, A. K. Gupta, G. Sheet, V. Chandrasekhar, H. W. Jang, and C. B. Eom, Pseudo-gap formation in the metallic state of La0.7Sr0.3MnO3 thin films, Appl. Phys. Lett. 93, 212503 (2008). [14] P. Dey, T. K. Nath, and A. Taraphder, Effect of substrate-induced strain on transport and magnetic properties of epitaxial La0.66Sr0.33MnO3 thin films, Appl. Phys. Lett. 91, 012511 (2007). [15] S. H. Kim, H. J. Sohn, Y. C. Joo, Y. W. Kim, T. H. Yim, H. Y. Lee, T. Kang, Effect of saccharin addition on the microstructure of electrodeposited Fe–36 wt.% Ni alloy, Surf. Coat. Tech. 199, 43 (2005). [16] P. Dey and T. K. Nath, Tunable room temperature low-field spin polarized tunneling magnetoresistance of La0.7Sr0.3MnO3 nanoparticles, Appl. Phys. Lett. 89, 163102 (2006). [17] D. Fiorani, Surface effect in magnetic nanoparticles, Springer (2005). [18] A. P. Ramirez, Colossal magnetoresistance, J. Phys. Conds. Mat. 9, 8171 (1997). [19] H. Y. Hwang, S. W. Cheong, N. P. Ong, and B. Batlogg, Spin-Polarized Intergrain Tunneling in La2/3Sr1/3MnO3, Phys. Rev. Lett. 77, 2041 (1996). [20] P. Raychaudhuri, T. K. Nath, A. K. Nigam and R. Pinto, A phenomenological model for magnetoresistance in granular polycrystalline colossal magnetoresistive materials: The role of spin polarized tunneling at the grain boundaries, J. Appl. Phys. 84, 2048 (1998). [21] J. Kondo, Resistance Minimum in Dilute Magnetic Alloys, Prog. Theo. Phys. 32, 37 (1964). [22] G. Bergmann, Weak localization in thin films: a time-of-flight experiment with conduction electrons, Phys. Rep. 107, 1 (1984). [23] P. A. Lee and T. V. Ramakrishnan, Disordered electronic systems, Rev. Mod. Phy. 57, 287 (1985). [24] P. Khatua, T. K. Nath, Mitali Banerjee, and A. K. Majumdar, Quantum interference effects and magnetic scattering in the electrical resistivity of Ni nanocrystallites in TiN matrix, Appl. Phys. Lett. 92, 193106 (2008). [25] A. H. Wilson, The Electrical Conductivity of the Transition Metal, Prog. R. Soc. London A 167, 580 (1938).

Chapter 6

Junction magnetoresistance study in

La0.7Sr0.3MnO3/SiO2/p-Si heterostructures


International journals 1. Electrical and magnetoelectronic properties of La0.7Sr0.3MnO3/SiO2/p‐Si heterostructure for spintronics application,

S. Chattopadhyay, P. Dey and T. K. Nath, Current Applied Physics (In press) doi:10.1016/j.cap.2011.02.009 2. On investigation of origin of junction magnetoresistance in La0.7Sr0.3MnO3/SiO2/p‐Si heterostructures, S.

Chattopadhyay and T. K. Nath, Journal of Physics D: Applied Physics (Communicated)

Conferences/Symposia 1. Investigation on La0.7Ca0.3MnO3/SiO2/n‐Si and La0.7Sr0.3MnO3/SiO2/p‐Si MOS like heterostructures for Spintronics by

S. Chattopadhyay, S. K. Giri and T. K. Nath, International Conference on Fundamental & Applications of Nanoscience and Technology (ICFANT) (2010).

2. Electrical properties of La0.7Sr0.3MnO3/SiO2/Si MOS structure by S. Chattopadhyay, P. Dey, T. K. Nath 53rd DAE Solid State Physics Symposium (2008)

3. I‐V characteristics of La0.7Sr0.3MnO3/SiO2/Si MOS structure by S. Chattopadhyay, P. Dey, T. K. Nath National Seminar on Advanced Nanomaterials and its Applications (2008)

Junction magnetoresistance in La0.7Sr0.3MnO3/SiO2/p‐Si heterostructures

Chapter 6

111

6.1. Introduction

After Datta-Das transistor [1], there is an enormous development of spintronic

devices in the area of microelectronics. Spintronics is a multidisciplinary field whose

central theme is the active manipulation of spin degrees of freedom in solid-state systems.

Generation of spin polarization usually means creating a non-equilibrium spin population

which can be achieved in several ways. For device applications electrical spin injection is

more desirable. In electrical spin injection a magnetic electrode is connected to the

sample. When the current drives spin-polarized electrons from the electrode to the

sample, non-equilibrium spin accumulates there. The rate of spin accumulation depends

on spin relaxation, the process of bringing the accumulated spin population back to

equilibrium. These properties of spin lead to open up new area of spintronic devices such

as magnetic sensors, magnetic devices, photosensitive devices etc [2-4]. Typical time

scales for spin relaxation in electronic systems are measured in nanoseconds, while the

range is from picoseconds to microseconds for electrons, which can be very useful for

high speed magneto-optic devices [5]. Spin detection, also a part of a generic spintronic

scheme, typically relies on sensing the changes in the signals caused by the presence of

non-equilibrium spin in the system. The common goal in many spintronic devices is to

maximize the spin detection sensitivity to the point that it detects not the spin itself, but

changes in the spin states.

Manganite-based heterojunctions such as half-metalic colossal magnetorsistance,

materials with extremely high degree of spin polarization [)()()()(

FF

FF

ENENENEN

S↓↑

↓↑

+

−= ] have

recently attracted lot of attention for their various fields of applications. Recently, a lot of

research studies have been focused on perovskite based oxide semiconductors or

semimetals heterojunctions with different semiconductors [6-8]. Generally the current-

voltage characteristic is guided by tunneling of conduction electrons between the

junctions though the presence of edge leakage makes the transport mechanism

complicated. As a result there is difficulty in the extraction of actual barrier height.

Lanthanum Strontium Manganese Oxide with composition La0.7Sr0.3MnO3 (LSMO), a

colossal magneto-resistive (CMR) ferromagnetic metal at room temperature belongs to

Chapter 6 Junction magnetoresistance study in La0.7Sr0.3MnO3/SiO2/p‐Si heterostructures

112

one of the perovskite type i.e., half-metallic highly spin-polarized ferromagnetic oxides,

hole doped manganites. The resistivity of the LSMO films can be tuned easily during

deposition and is highly stable under further thermal treatment. Moreover, the resistivity

of the materials can also be tuned by applying external magnetic field [9]. So using

LSMO as an electrode in Si/ SiO2/ LSMO heterostructure the I-V and C-V characteristics

can be modulated by applying magnetic field externally and it may open up possibilities

in various technological applications, e.g., magnetic sensors, magnetic memories or other

spin electronic applications. However, the origin of appearing positive junction MR of

these perovskite based heterojunctions is still not understood clearly. Some reports have

concluded that the increase of barrier height due to Zeeman splitting of bands causes the

positive MR at the junction. But, the exact reason is still far from well established.

In this chapter, the junction current density-voltage (J-V) characteristics of

La0.7Sr0.3MnO3 (LSMO)/SiO2/Si junction with and without external magnetic field have

been studied explicitly. The applied external magnetic fields up to 1 T and 8 T have been

exerted on all the heterojunctions using a high precision electromagnet. The constant

parameters of J-V characteristics at room temperature and all other temperature have

been estimated by non-linear least square (χ2- minimization) curve fitting method. The

estimated parameters conclusively show that the series resistance of the junction plays

dominant role to make junction MR positive rather than the effective barrier height. The

effect of leakage currents on junction MR for such heterostructures has also been

established.


The details of sample preparations have been discussed in details in chapter 5

section 5.2. The three best films which show higher value of negative magnetoresistance,

(Film magnetoresistance) have been chosen for heterojunction study. The heterojunctions

and its growth structures and conditions have been summarized in Table 6.1. The

variation of the oxidation time at oxidation temperature 800 °C under oxygen atmosphere

of cleaned high conducting p-Si substrate forms different thermal oxide layer over p-Si

with different thicknesses. The insulating layer between p-Si and La0.7Sr0.3MnO3 (LSMO)

controls the current transport through the heterojunctions.


Chapter 6

113

Table 6.1. The heterojunction structures and its growth conditions

Heterojunction name

Structure of heterojunction

Growth conditions for heterojunction

Film properties

Sample-1 p-Si/SiO2 (native oxide)/

La0.7Sr0.3MnO3 Substrate p-Si; No oxidation ((thickness of SiO2 ~9 Å) La0.7Sr0.3MnO3 have been grown at 800 °C substrate temperature under 0.5 mbar O2 atmosphere

La0.7Sr0.3MnO3

shows rod like structure with -ve magneto-resistance of film

Sample-2 p-Si/SiO2(thermal oxide)/ La0.7Sr0.3MnO3

Substrate p-Si; Oxidation has been carried out for 30 min at 800 °C under oxygen pressure 0.5 mbar (thickness of SiO2 ~40 Å). Then the La0.7Sr0.3MnO3 film have been grown at 800°C substrate temperature under 0.5 mbar O2 atmosphere

La0.7Sr0.3MnO3

shows rod like structure with –ve magneto-resistance film

Sample-3 p-Si/SiO2(thermal oxide)/ La0.7Sr0.3MnO3

Substrate p-Si; Oxidation has been carried out for 45 min at 800 °C under oxygen pressure 0.5 mbar (thickness of SiO2 ~45 Å). Then the La0.7Sr0.3MnO3 film

have been grown at 800 °C substrate temperature under 0.5 mbar O2 atmosphere

La0.7Sr0.3MnO3

shows rod like structure with –ve magneto-resistance in the form of thin film.

Electrical contacts were made with high purity Ag on LSMO film and pure Al contact

with p-Si as ohmic contacts. The temperature and magnetic field dependence I-V

characteristics were measured using a source meter (Keithley 2612), current source

(Keithley 6221), temperature controller (Lakeshore, model-331), high precision

electromagnet (Polytronic, Model HEM 100) and a variable temperature cryostat (Janis,

USA). A cryogen free 8 T superconducting magnet with VTI system down to temperature


114

2 K (Cryogenics, U.K.) has been employed for high field and low temperature transport

measurements.

6.3 Results and discussion

6.3.1 Structural properties

The x-ray diffraction pattern of LSMO film (sample-3) deposited on (100) p-Si

substrate using Cu-Kα radiation has been shown in Fig. 6.1. The multipeaks from

different crystallographic planes of LSMO sample reveal the non-epitaxial nature of the

LSMO film on SiO2/Si layer as discussed in chaper-5. In Fig. 6.1(b) and 6.1(c) the

FESEM micrograph of the LSMO film (sample-3) and the cross sectional FESEM image

of the heterojunction have been shown. The FESEM micrograph clearly shows the

uniformly grown smooth LSMO film with good coverage on Si/SiO2 layer. The cross-

sectional FESEM picture reveals the nano-rod like growth of the film with film thickness

~ 600 nm.

6.3.2. Electrical properties of LSMO/SiO2/p-Si hereostructure without applied magnetic

field

6.3.2.1. Current-Voltage study using diode characteristics

Figure 6.2(a) shows the non-linear current density-voltage (J-V) characteristics of

a typical LSMO/SiO2/Si structured heterojunction measured at room temperature. It can

(a)

(c)

(b)

(c)

(a)

Fig. 6.1. (a) XRD pattern of LSMO/SiO2/p-Si hereostructure. Inset shows the XRD pattern of LSMO film. (b) and (c) FESEM micrograph of LSMO film and crosssectional view of the heterostructure.


Chapter 6

115

be seen that the junctions exhibit diode-like behavior. The J-V characteristics of the films

are found to depend strongly on the interfacial oxides layer. The ideality factor η is much

greater than 2 at room temperature. At low forward voltage the current increased

exponentially which has been usually observed in diodes and attributed to recombination,

tunneling mechanism. At moderate junction voltage the J-V deviated from ideal

thermionic emission and behaves as J~V2 relation which is attributed to space charge

limited current (SCLC) conduction. Applied voltage lower than the turn on voltage shows

ohmic behavior. All the regions have been distinctly shown in log J - log V plot of Fig.

6.2(b). The junction J-V characteristics are described as [10],

⎟⎟⎠

⎞⎜⎜⎝

⎛ −−=

kTIRVVe

JJ s

η)(

exp 00 (6.1)

where, J0 is the reverse saturation current density. V0, η, and Rs are turn on voltage,

ideality factor and junction series resistance respectively. The fitted values of the

parameters for different samples have been plotted in Fig. 6.2(c) and Fig. 6.2(d),

respectively. The decrease of ideality factor and reverse saturation current density [in Fig.

6.2(c)] and increase of turn on voltage and series resistance [in Fig. 6.2(d)] with

increasing the oxidation time illustrates that the leakage current and defect induced

tunneling through the junction decreases effectively with the enhancement of oxidation

time (thickness of the intermediate SiO2 oxide layer). As the native oxide layer contains

inherently high density of oxide defect states in it, the leakage current becomes large at

room temperature that enhances the junction forward current and causes lower series

resistance through the junction. When the oxide layer is grown over the Si substrate

thermally in O2 atmosphere the oxide defects decreases as well as the thickness of the

oxide layer increases. So, both the leakage current and tunneling probability through the

oxide interface decreases. It causes enhancement of series resistance from sample-1 to

sample-3.


116

The J-V characteristics for sample-3 at several temperatures down to 77 K have

been shown in Fig. 6.3(a). It shows that the current decreases with decreasing

temperature. The best fit parameters (η, J0, Rs and V0) for all three samples have been

evaluated for all six temperatures down to 77 K using Eq. (6.1) employing a χ2

minimization technique. All these temperature dependent fit parameters are shown in Fig.

6.3(b) and 6.3(c). The strong dependency of current density and all four parameters with

temperature indicates that not only the thermionic emission is occurring at the junction

but other types of transport mechanisms are also present for all the heterojunctions. The

strong dependency of ideality factor with temperature implies that the tunneling

mechanism is one of the dominating current transport mechanisms through the junction.

-6 -4 -2 0 2 4 6-0.15

-0.10

-0.05

0.00

0.05

0.10

0.15

0.20

sample-1 sample-2 sample-3

C

urre

nt d

ensi

ty (m

A/c

m2 )

Voltage (V)

300 K

1 10

1

10

100

Cur

rent

den

sity

(μA

/cm

2 )

Voltage (V)

Ohmic region

Tunneling region

SCLC region

Sample-2

300 K

0 10 20 30 40 508

12

16

20

24

28

32

η J0


η

300 K

1.0

1.5

2.0

2.5

3.0

3.5

4.0

J0 (μA/cm

2)

0 10 20 30 40 500

10

20

30

40

50

60

70

80

RS

V0


Seri

es R

esis

tanc

eRS(

kΩ−c

m2 )

300 K

0.50

0.52

0.54

0.56

0.58

0.60

0.62

Turn on voltage V

0 (V)

(a) (b)

(c) (d)

Fig.6.2: (a) J-V characteristics of LSMO/SiO2/p-Si heterojunction for three different samples at room temperature. (b) log J-log V plot of sample-2 to understand the current transport process at room temperature. (c) and (d) are evaluated fit parameters at room temperature. The ideality factors, reverse saturation current density plot for three different samples have been shown in (c), and series resistance and turn-on voltage have been shown in (d).


Chapter 6

117

The ln(J0/T2) vs. 1000/T plot shown in Fig. 6.3(d) provides the information about the

effective barrier height and Richardson constant for all the three samples.

Table 6.2. Effective barrier height and Richardson constants for different samples

Sample name Effective barrier height (eV) Richardson constant (AK-2cm-2)

Sample-1 0.021 7.05 × 10-9

Sample-2 0.027 1.7 × 10-8

Sample-3 0.030 7.62 ×10-9

The evaluated values have been summarized in the Table 6.2. The slight difference in the

value of Richardson constants may be due to different effective masses in the films. The

0 2 4 6 8 100

20

40

60 77 K 120 K 150 K 200 K 250 K 300 K

Cur

rent

den

sity

(μA

/cm

2 )

Voltage (V)

Sample-3

(a)

50 100 150 200 250 3000

100

200

300

400

500

sam

ple-

3 sample-2

sample-1

sample-3

sample-2

sample-1

Temperature (K)

Seri

es r

esis

tanc

e (kΩ

-cm

2 )

0

2

4

6 RS V0

Turn on voltage (V

0 ) (V)(c)

2 4 6 8 10-19.6

-19.2

-18.8

-18.4

-18.0

-17.6

-17.2

-16.8

-16.4

Sample-2

Sample-1

ln(J

0/T

2 )

1000/T (K-1)

Sample-3

(d)

100 150 200 250 300

10

20

30

40

50

60

70

80

sample-3

sample-2

sample-1

sample-3

sample-2

sample-1

Temperature (K)η

0

1

2

3

4 η J0 J

0 (μA/cm

2)

(b)

Fig. 6.3. (a) Temperature dependent J-V characteristics of sample-3 without applied magnetic field. (b) and (c) are the plot of evaluated parameters with temperature. The ideality factor and reverse saturation current density plot with temperature for three different samples have been shown in (b), and series resistance and turn-on voltage with temperature have been shown in (c). (d) ln(I0/T2) vs. 1000/T plot to determine the effective barrier height.


118

barrier height is much lower than expected because of high leakage current through the

junction.

6.3.2.2. Tunneling Characteristics

The I-V properties reveal that tunneling is occurring at all temperatures along

with other mechanisms through our heterojunction as discussed earlier. Now it is

necessary to find out the possible tunneling mechanism through the heterojunctions. The

observed field stimulated emission and capture have been discussed in terms of Fowler-

Nordheim tunneling [11], Poole-Frenkel effect [12–14], phonon assisted tunneling [15],

and a combination of both phenomena [16-20]. Fowler-Nordheim tunneling is the process

whereby electrons tunnel through a barrier in the presence of a high electric field. This

quantum mechanical tunneling process is an important mechanism for thin barriers

similar to those in metal-semiconduictor junctions on highly-doped semiconductors.

Fowler-Nordheim (F-N) tunneling current into the SiO2 conduction-band through a

triangular barrier at high field is given by,

)/exp(2 VBAVJ FN −= (6.2)

where, Bh

qAφπ8

3= and

qhm

B BFN

3)2(8 2/32/1 φπ

= . Here q is the electronic charge, Bφ and mFN

are the effective tunneling barrier height and effective mass of electron for F-N tunneling.

V is the applied electric field across the thin oxide layer. The high field Fowler-Nordheim

(F-N) plot [ln(JFN/V2) vs 1/V] of the heterojunction with native (sample-1) and thermal

oxide intermediate layer (sample-3) at different temperatures have been shown in Fig.

6.4(a) and 6.4(b). The high field Fowler-Nordheim tunneling should be temperature

independent. But our observed [ln(JFN/V2) vs 1/V] is strongly temperature dependent and

hence we can conclude that the field dependent tunneling is mainly temperature

dependent Frenkel-Poole type emission.

The well-known Frenkel-Poole effect describes the increase of the thermal

emission rate of carriers in an external electric field due to the lowering of the barrier

associated with their Coulomb potential. It is a classical mechanism in which the electron

is thermally emitted over the top of a potential barrier which has been lowered by the

presence of an electric field.


Chapter 6

119

Frenkel-Poole emission refers to electric-field-enhanced thermal emission from a

trap state into a band of electronic states of insulator conduction band. The current

density associated with Frenkel-Poole emission is given by [21],

⎥⎥⎦

⎤

⎢⎢⎣

⎡ −−=

kTqVq

CVJ sBFP

επεφ 0/(exp (6.3)

Where, Bφ is the barrier height for electron emission from the trap state, sε is the relative

dielectric permittivity at high frequency, T is temperature, 0ε is the permittivity of free

space. As the electrons emitted from the trap states and can not polarize the surrounding

atoms, the relevant dielectric constant is that at high frequency, rather than the static

dielectric constant [15]. The high field ln(JFP/V) vs. V plot shown in Fig. 6.5(a) and

Fig. 6.5(b) are linear for all temperatures and it implies that the temperature dependent

Frenkel-Poole emission is the dominating current transport mechanism in non-epitaxial

LSMO/SiO2/p-Si heterostructures.

0.1 0.2 0.3 0.4 0.5

-14

-12

-10

-8

-6

-4300 K

250 K

200 K

ln(J

FN/V

2 )

1/V (V-1)

Sample-1 150 K

(a)

0.1 0.2 0.3 0.4 0.5

-14

-12

-10

-8

-6

300 K

250 K

200 K

ln(J

FN/V

2 )

1/V (V-1)

150 KSample-3 (b)

Fig.6.4. ln(JFN/V2) vs 1/V plot at different temperatures for (a) sample-1 and (b) sample-3.


120

These plots show linear nature in higher field range confirms the presence of temperature

dependent Frenkel-Poole emission through the heterostructures.

6.3.3. Electrical properties of LSMO/SiO2/p-Si hereostructure with applied magnetic field

The J-V behaviors of sample-3 under very high magnetic field up to 8 T have

been shown in Fig. 6.6(a) and Fig. 6.6(b) at 300 K and 120 K, respectively.

Fig. 6.5. ln(JFP/V) vs. plot for V1/2 at different temperatures for (a) sample-1 and (b) sample-3

0.5 1.0 1.5 2.0 2.5 3.0

-14

-12

-10

-8

-6

-4

300 K

250 K

200 K

√V (V1/2)

ln(J

FP/V

)

150 K

(b)

Sample-3

0.5 1.0 1.5 2.0 2.5 3.0 3.5

-14

-12

-10

-8

-6

-4

300 K

250 K

200 K

ln(J

FP/V

)

√V (V1/2)

150 KSample-1

(a)

Fig. 6.6. The J-V characteristics at 300 K and 120 K in (a) and (b), respectively at different applied magnetic field; the %JMR at 300 K and 120 K in (c) and (d), respectively.

0 1 2 3 4 5 6

0.000.010.020.030.040.050.060.07

0 T 1 T 2 T 4 T 6 T 8 T

Cur

rent

den

sity

(mA

/cm

2 )

Voltage (V)

Sample-3T = 300 K

(a)

0 2 4 6 8 10

0.000

0.002

0.004

0.006

0.008

0.010

Cur

rent

den

sity

(mA

/cm

2 )

Voltage (V)

0 T 1 T 2 T 4 T 6 T 8 T

Sample-3T =120 K

(b)

-8 -6 -4 -2 0 2 4 6 8

0

10

20

30

40

50

60

%JM

R

Magnetic field (T)

Sample-3120 K

(d)

-8 -6 -4 -2 0 2 4 6 8

0

4

8

12

16

20

% J

MR

Magnetic field (T)

Sample-3300 K

(c)


Chapter 6

121

Corresponding junction magnetoresistance have also been shown in Figs. 6.6(c)

and (d), respectively. The change is almost linear up to 1 T magnetic field. After applying

higher magnetic field beyond 1 T, the junction magnetoresistance starts saturating and

applying very high magnetic field (> 4 T), junction magnetoresistance starts to decrease

slightly.

6.3.3.1. Current-Voltage properties under magnetic field using diode characteristics

Figure 6.7 (a) shows the room temperature J-V nature with and without external

magnetic field of 1 T for all three samples.

The J-V curve shows reasonably good sensitivity under magnetic field at room

temperature as shown in Fig. 6.7(a). The four parameters in Eq. (6.1) are modified under

0 10 20 30 40 5012141618202224262830

ηm J0m


η mT=300 K

0.014

0.016

0.018

0.020

(b)

J0m (m

A/cm

2)

0 10 20 30 40 500

20

40

60

80

100

RSm

V0m


RSm

(kΩ

-cm

2 )

1.0

1.1

1.2

1.3

1.4

1.5

(c)

V0m (V

)

T =300 K

0 1 2 3 4 5 60.00

0.06

0.12

0.18 sample-1 at 0T sample-1 at 1T sample-2 at 0T sample-2 at 1T sample-3 at 0T sample-3 at 1T

curr

ent d

ensi

ty (m

A/c

m2 )

Voltage (V)

(a)

Fig. 6.7(a) J-V characteristics of LSMO/SiO2/p-Si heterojunction for three different samples at room temperature with and without applied 1 T magnetic field. (b) and (c) are evaluated parameters at room temperature. The ideality factor and reverse saturation current density plot for three different samples under 1 T magnetic field have been shown in (b), and series resistance and turn-on voltage have been shown in (c) at room temperature. (d) The measured %junction MR plot with applied magnetic field for all three samples at room temperature.

-0.8 -0.4 0.0 0.4 0.802468

10121416

(d)

Sample-1 Sample-2 Sample-3

% J

MR

Applied magnetic field (T)

T=300 K


122

magnetic field (J0, η, Rs and V0 are now replaced by J0m, ηm, Rsm and V0m, respectively)

and the current density is found to decrease with magnetic field at higher bias potential.

The reverse saturation current density (J0m) and ideality factor (ηm) evaluated from fitting

to room temperature J-V curve under magnetic field for different oxide thickness have

been plotted in Fig. 6.7(b) and series resistance (Rsm) and turn-on voltage (V0m) have

been shown in Fig. 6.7(c). The current has been measured by keeping fixed forward

junction voltage at 4.8 V for varying magnetic field up to 1 T at room temperature. It

shows the change in junction MR is about 7, 9 and 16% at 4.8 V for sample-1, sample-2

and sample-3, respectively. The definition used here for the junction

%100)0()0()1(×⎟⎟⎠

⎞⎜⎜⎝

⎛⎟⎟⎠

⎞⎜⎜⎝

⎛−=

=== THJ

applied

THJ

applied

THJ

applied

IV

IV

IV

MR , where IJ(H=1T) is the current at 1 T applied

magnetic field. The plots of % of junction MR with applied magnetic field at a bias

voltage of 4.8 V for all three samples have been shown in Fig. 6.7(d).

The temperature dependent J-V properties without and with 1 T magnetic field of

sample-3 have been shown in Fig. 6.8(a) for three different temperatures. It shows the

positive junction MR properties for all the temperatures and the magnitude of junction

MR is also temperature dependent. The fit parameters under magnetic field for different

samples have been plotted in Fig. 6.8(b) and Fig. 6.8(c) respectively. The dependency of

all four parameters with temperature show the same trend as the parameters evaluated

from J-V without magnetic field shows. The ln(J0m/T2) vs. 1000/T plot shown in Fig.

6.8(d) provides the information about the effective barrier height and Richardson constant

for all the three heterojunctions under magnetic field. The estimated parameters have also

been listed in the Table 6.3. The evaluated effective barrier heights are not much

increased under magnetic field. But the other parameters, mainly the series resistances are

increased significantly under magnetic field and it cause the positive MR at the

heterojunctions at all operating temperatures.


Chapter 6

123

Table 6.3. Effective barrier height and Richardson constants for three different samples under 1 T magnetic field

Sample name Effective barrier height (eV) Richardson constant AK-2cm-2

Sample-1 0.020 2.73 × 10-8

Sample-2 0.029 4.22 × 10-8

Sample-3 0.046 9.21 × 10-8

The comparative study of barrier height and series resistance with and without magnetic

field has been listed in Table 6.4.

0 2 4 6 8 100.000.010.02

0.030.040.05

0.060.07

(a)

120 K at 0T 120 K at 1T 250 K at 0T 250 K at 1T 300 K at 0T 300 K at 1T

curr

ent d

enst

y (m

A/c

m2 )

Voltage (V)

Sample-3

120 160 200 240 280 320

0

100

200

300

400

500

600

(c)

Sample-3

Sample-2

Sample-1

Sample-3Sample-2

RSm

V0m

Temperature (K)

Rsm

(KΩ−c

m2 )

Sample-1

2

4

6

V0m

(V)

3 4 5 6 7 8 9-19.0-18.5-18.0-17.5-17.0-16.5-16.0-15.5-15.0

(d)ln(J

0m/T

2 )

1000/T (K-1)

sample-3

sample-2

sample-1

120 160 200 240 280 320

20

40

60

80

100

Sam

ple-1

Sample

-3

Sample-3

Sample-2

ηm J0m

Temperature (K)

η m

Sample-1

0.000

0.004

0.008

0.012

0.016

0.020

(b)

J0m (m

A/cm

-2)

Fig. 6.8. (a) Temperature dependent J-V characteristics of sample-3 with and without 1 T applied magnetic field. (b) and (c) are the plot of evaluated parameters with temperature. The ideality factor and reverse saturation current density plot with temperature for three different samples under 1 T magnetic field have been shown in (b), and series resistance and turn-on voltage with temperature under 1 T magnetic field have been shown in (c). (d) ln(I0/T2) vs. 1000/T plot to determine the effective barrier height under magnetic field.


124

Table 6.4. Comparative study of effective barrier height and room temperature series resistance with and without applied 1 T magnetic field.

Applied Magnetic

field

Sample name Effective barrier height

(eV)

Series resistance

(kΩ-cm2)

0 T Sample-1 0.021 7.2

Sample-2 0.027 35.42

Sample-3 0.030 74.86

1 T Sample-1 0.020 7.8

Sample-2 0.029 39.34

Sample-3 0.046 89.03

It is clear that the evaluated effective barrier heights are not much increased under

magnetic field. But the other parameters, mainly the series resistances are increased much

under magnetic field and cause the positive MR at the heterojunctions at all operating

temperatures.

6.3.3.2. Tunneling Characteristics under1 T applied magnetic field

As discussed earlier, the dominating tunneling mechanism through the

heretojunctions fits very well with temperature dependent Frenkel-Poole emission. The

ln(JFP/V) vs. V plot for J-V under 1 T applied magnetic field has also been shown in

Fig 6.9(a) and Fig. 6.9(b) for both sample-1 and sample-3, respectively.

In the high biasing field region ln(JFP/V) vs. V graphs have been fitted linearly for both

the sample-1 and sample-3. The intercept at y-axis have been plotted with 1000/T in Fig.

Fig.6.9. (a) ln(JFP/V) vs. plot for V1/2 at different temperatures for (a) sample-1 and (b) sample-3 under 1 T applied magnetic field.

0.5 1.0 1.5 2.0 2.5 3.0 3.5

-14

-12

-10

-8

-6

-4 300 K 250 K

200 K

ln(J

FP/V

)

√V (V1/2)

Sample-1H= 1T

150 K

(a)

0.5 1.0 1.5 2.0 2.5 3.0 3.5-15-14-13-12-11-10-9-8-7-6-5

300 K

250 K

200 K150 K

ln(J

FP/V

)

√V (V1/2)

Sample-3H = 1 T

(b)


Chapter 6

125

6.10 (a) and (b) for sample-1 and sample-3 without and with 1 T applied magnetic field,

respectively. The slope of the plots gives the effective tunneling barrier height.

The evaluated effective tunneling barrier height with and without magnetic field for three

samples have been listed in Table 6.5.

Table 6.5: Evaluated effective tunneling barrier height with and without 1 T applied magnetic field for different samples

Sample Applied magnetic field

(Tesla)

Effective tunneling barrier

height (meV)

Sample-1 0 179.72

1 180.37

Sample-2 0 193.54

1 194.11

Sample-3 0 211.21

1 212.70

6.3.4. Junction magnetoresistance properties study

The % of JMR for the LSMO/SiO2/p-Si heterojunction has been calculated from

the relation, %100×−

=s

ssm

RRR

JMR . The room temperature magnetic field dependent

2.5 3.0 3.5 4.0 4.5 5.0 5.5 6.0 6.5 7.0 7.5-16

-14

-12

-10

-8

-6

Sample-3

Inte

rcep

t at y

-axi

s

1000/T (K-1)

Sample-1

(a)

H = 0 T

2.5 3.0 3.5 4.0 4.5 5.0 5.5 6.0 6.5 7.0 7.5-16

-14

-12

-10

-8

-6 Sample-3Sample-1

Inte

rcep

t at y

-axi

s

1000/T (K-1)

(b)

H = 1 T

Fig. 6.10. The intercept at y-axis evaluated from Fig. 6.5 and Fig. 6.9 with 1000/T plot for sample-1 and sample-3 under (a) 0 T and (b) 1 T applied magnetic field to determine the effective tunneling barrier height


126

junction magnetoresistance (JMR) behavior in the low magnetic field region (up to 1 T)

of sample-3 have been analyzed using a simple empirical relation as [22],

βαHJMR =% (6.4)

where, α and β are coefficients. The best fit values of α and β are and are found to be

0.31 and 0.73, respectively, employing a non-linear least square fitting χ2 minimization

technique. The coefficient β is found to be smaller than one at room temperature showing

non-linear magnetic field dependence of positive MR of the junction. The high

dependency of JMR with interfacial oxide layer implies that the leakage current also

plays an important role in JMR property of such p-i-p type heterojunctions. As discussed

earlier, the increase of oxidation time causes the increase of crystalline SiO2 oxide layers

and thereby causes decrease of oxide defects. The less oxide defect causes less leakage

current through the junction. As defect related leakage currents decreases for sample-3,

the value of the parameters (coefficients) enhances (modifies) and causes higher % of

junction MR at room temperature [Fig. 6.7(d)].

Table 6.6. Evaluated α and β parameters

Heterojunctions α β (T-β)

Sample-1 8.42 0.85

Sample-2 9.89 0.85

Sample-3 16.32 0.80

The increase of junction series resistance can be explained by the theoretical

model of spin tunneling in ferromagnetic to paramagnetic junctions [23]. It is interesting

that the existence of JMR is positive when the dominating current transports mechanism

is tunneling. So, it can be concluded that the appearance of positive JMR is due to

tunneling of electron through the heterojunction. Considering eg electron tunnels from

ferromagnetic LSMO to paramagnetic p-Si through insulating SiO2 layer, it can be


Chapter 6

127

considered that there exist two channels (spin up I↑ and spin down I↓) at conduction band

for eg electrons to tunnel as shown in Fig. 6.11.

The ratio I↑/I↓ is proportional to the spin polarization of LSMO (Fig. 6.11). At room

temperature the ratio is approximately 1 due to the lower spin polarization of LSMO. The

both channels have allowed states at Fermi energy level and the both channels act in

tunneling mechanism at room temperature in absence of magnetic field. If the external

magnetic field is applied, the disordered spins are suppressed and the I↑/I↓ ratio increases.

One channel becomes inactive due to applied magnetic field and at Fermi energy level

only the other channel will have allowed states for tunneling at room temperature and low

temperatures. This reduces the current at a particular biasing field in presence of 1 T

magnetic field as observed in the J-V characteristics of the heterojunction. It most likely

causes enhancement of junction resistance under applied magnetic field.

The slight decrease of turn-on voltage with applied magnetic field shown in Fig.

6.3 and Fig. 6.8 implies that there is a certain voltage region in which the junction MR is

negative. This voltage region lies between the turn-on voltage and saturation voltage for

corresponding temperatures. The % of JMR plot with temperature has been shown in Fig.

6.12. The % JMR for all three samples decreases sharply with increasing the temperature.

Fig. 6.11. Schematic diagram of spin tunneling from magnetic LSMO to non magnetic Si with and without applied magnetic field.


128

With increasing of thermal energy it is more likely that the spin polarization decreases

and hence the change of junction resistance decreases. It causes the decrease of % JMR

with increasing temperature.

The change of junction MR with different forward bias voltages for all the three

samples at room temperature and for different temperatures for sample-3 have been

shown in Fig. 6.13(a) and Fig. 6.13(b), respectively.

The positive junction MR occurs at higher voltages and increases almost exponentially

with increasing forward bias voltages for all samples at all temperatures. In lower bias

Fig. 6.12. The temperature dependent junction MR plots for all three samples; the JMR has been calculated from the values of series resistances with and without 1 T magnetic field which are evaluated from the fitting of J-V curves.

120 160 200 240 280 3205

101520253035404550


% J

MR

Temperature (K)

2 3 4 5 60

4

8

12

16

20

24


% J

MR

Applied voltage (V)

T = 300 K

1 10

10

20

30

40

50

120 K 250 K 300 K

% J

MR

Applied Voltage (V)

Sample 3

Fig. 6.13. (a) The % of junction MR plot with applied bias voltage for all three samples. (b) The same plot for different temperature for sample-3.


Chapter 6

129

voltage regions (voltage lies between turn-on voltage and saturation voltage) the junction

shows either no such significant change in junction MR or negative junction MR for

different samples. The negative junction MR in this region becomes prominent at low

temperatures rather than room temperatures. It may be due to the effect of leakage current

through the junction. The leakage current is higher at room temperatures and there is no

such significant effect of magnetic field on leakage current which has been found through

this study. The little change of junction current may be compensated with leakage current

and overall current change can vanish at room temperature. When temperature decreases

the leakage current also decreases. The current change through the junction becomes

prominent at low temperatures.

6.4. Summary

In summary, we have successfully grown the LSMO/SiO2/p-Si heterostructure by

pulsed laser deposition technique and investigated its junction properties in details. We

have varied the SiO2 barrier layers thickness by varying oxidation times on cleaned Si

substrates. The various oxidation conditions are responsible for generating defects in

oxide layer which most likely play important role in carrier transport through the

junction. The dominating current transport mechanism through the junction is

temperature dependent Frenkel-Poole type emission. We have thoroughly investigated

the dependency of positive junction MR with applied biasing potentials and temperatures.

We have also found that the junction MR also depends on the defects induced leakage

current through the junction. The junction MR increases with increasing the thickness of

the SiO2 interfacial layer. It can be concluded that the leakage current can damage the

magnetic sensitivity (decrease of junction MR) of such devices. We have also estimated

the junction parameters such as ideality factor, turn-on voltages, barrier heights, series

resistances etc. It is found that the ideality factor, barrier height and series resistance

increase and turn-on voltages slightly decreases with the increase of external applied

magnetic field at all operating temperatures for all three samples. These results reveal that

the junction MR is positive at higher potentials and shows little negative value at lower

biasing voltage regions. Highest junction MR for the sample-3 has been found to be

~56% at 120 K and ~17% at 300 K at an applied bias voltage of 2 T. The thorough


130

analysis of junction J-V characteristics also leads to conclude that the arising of positive

MR at the junction is mainly due to the change of series resistances under magnetic field

rather than the change of effective barrier heights. We have also tried to establish

possible mechanism of the observed experimental results from our oxide thickness layer

dependent heterojunction employing the possible model of FM semiconductor/NM

semiconductor tunneling junction.

References

[1] S. Datta and B. Das, Electronic analog of the electro‐optic modulator, Appl. Phys. Lett. 56, 665 (1990). [2] J. Fontcuberta, L. Balcells, M. Bibes, J. Navarro, C. Frontera, J. Santiso, J. Fraxedas, B. Martínez, S. Nadolski, M. Wojcik, E. Jedryka and M J, Casanove, Magnetoresistive oxides: new developments and applications, J. Magn. Magn. Mat. 242, 98 (2002). [3] Y. Ishii, H. Yamada, H. Sato, H. Akoh, M. Kawasaki, Y. Tokura and Y. Tokura, Perovskite manganite magnetic tunnel junctions with enhanced coercivity contrast, Appl. Phys. Lett. 87, 022509 (2005). [4] R. S. Popovic, Not-plate-like Hall magnetic sensors and their applications, Sens. Act. A 85, 9 (2000). [5] M. Johnson, Magnetoelectronics, Academic Press, Elsevier, London,(2004). [6] J. Xing, K. Zhao , G. Z. Liu, M. He, K. J. Jin and H. B. Lu, Enhancement of photovoltaic effect in La0.9Sr0.1MnO3/Si heterojunction by side illumination, J. Phys. D: Appl. Phys. 40 5892 (2007). [7] K. Zhao, K. J. Jin, H. Lu, Y. Huang, Q. Zhou, M. He, Z. Chen, Y. Zhou and G. Yang, Transient lateral photovoltaic effect in p-n heterojunctions of La0.7Sr0.3MnO3 and Si, Appl. Phys. Lett. 88, 141914 (2006). [8] Lord K, Hunter D, Williams T M and Pradhan A K, Photocarrier injection effect and p-n junction characteristics of La0.7Sr0.3MnO3/ZnO and Si heterostructures, Appl. Phys. Lett. 89, 052116 (2006). [9] R. L. Zhang, W. H. Song, Y. Q. Ma, J. Yang, B. C. Zhao, G. H. Zheng, Z. G. Sheng, W. J. Lu and Y. P. Sun, The response for magnetic field, current and photo irradiation of charge-ordering LaSr2Mn2O7 thin film, J. Phys. D: Appl. Phys. 39 621 (2006). [10] S. J. May and B. W. Wessels, High-field magnetoresistance in p-(In,Mn)As/n-InAs heterojunctions, Appl. Phys. Lett. 88, 072105 (2006). [11] G. Chakraborty, S. Chattopadhyay, C. K. Sarkar, and C. Pramanik, Tunneling current at the interface of silicon and silicon dioxide partly embedded with silicon nanocrystals in metal oxide semiconductor structures, J. Appl. Phys. 101, 024315 (2007). [12] S. D. Ganichev, E. Ziemann, W. Prettl, I. N. Yassievich, A. A. Istratov and E. R. Weber, Distinction between the Poole-Frenkel and tunneling models of electric-field-stimulated carrier emission from deep levels in semiconductors, Phys. Rev. B 61, 10361 (2000). [13] G. Vincent, A. Chantre, and D. Bois, Electric field effect on the thermal emission of traps in semiconductor junctions, J. Appl. Phys. 50, 5484 (1979).


Chapter 6

131

[14] W. R. Buchwald, N. M. Johnson, Revised role for the Poole-frenkel effect in deep-level characterization, J. Appl. Phys. 64, 958 (1988). [15] L.Tsybeskov, G. F. Grom, P. M. Fauchet, J. P. McCaffrey, J. M. Baribeau, G. I. Sproule, and D. J. Lockwood , Phonon-assisted tunneling and interface quality in nanocrystalline Si/amorphous SiO2 superlattices, Appl. Phys. Lett. 75, 2265 (1999). [16] G. Vincent, A. Chantre, and D. Bois, Electric field effect on the thermal emission of traps in semiconductor junctions, J. Appl. Phys. 50, 5484 (1979). [17] F. D. Auret, S. A. Goodman, and W. E. Meyer, Electric-field-enhanced emission from radiation-induced hole traps in p-GaAs, Semicond. Sci. Technol. 10, 1376 (1995). [18] A. Ilie and B. Equer, Field-enhanced generation in hydrogenated amorphous silicon, Phys. Rev. B 57, 15349 (1998). [19] W. R. Buchwald and N. M. Johnson, Revised role for the Poole-frenkel effect in deep-level characterization, J. Appl. Phys. 64, 958 (1988). [20] P. A. Martin, B. G. Streetman and K. Hess, Electric field enhanced emission from non-Coulombic traps in Semiconductors, J. Appl. Phys. 52, 7409 (1981). [21] H. Zhang, E. J. Miller and E. T. Yua, Analysis of leakage current mechanisms in Schottky contacts to GaN and Al0.25Ga0.75N/GaN grown by molecular-beam epitaxy, J. Appl. Phys. 99, 023703 (2006). [22] Z. G. Sheng, W. H. Song, Y. P. Sun, J. R. Sun and B. G. Shen, Crossover from negative to positive magnetoresistance in La0.7Ce0.3MnO3-SrTiO3-Nb heterojunctions, Appl. Phys. Lett. 87, 032501 (2005). [23] I. Žutić, J. Fabian, and S. Das Sarma, Spintronics: Fundamentals and applications, Rev. Mod. Phys. 76, 323 (2004).

Chapter 7

Electronic-and magneto transport of

La0.7Sr0.3MnO3/ZnO, La0.7Sr0.3MnO3/Zn(Fe)O and

La0.7Sr0.3MnO3/Zn(Fe,Al)O heterostructures


International journals 1. Enhanced temperature dependent junction magnetoresistance in the heterojunctions with La0.7Sr0.3MnO3 and iron

doped ZnO carrier induced dilute magnetic semiconductors by S. Chattopadhyay, J. Panda, T. K. Nath, Journal of Applied Physics. (Communicated)

Conferences/Symposia 1. Temperature dependent junction magnetoresistance behavior of LSMO/Zn(Fe,Al)O heterojunction for spintronics by

J. Panda, S. Chattopadhyay and T. K. Nath, 55th DAE Solid State Physics Symposium 2010 (2010).

Electronic‐and magneto transport of La0.7Sr0.3MnO3/ZnO, La0.7Sr0.3MnO3/Zn(Fe)O and La0.7Sr0.3MnO3/Zn(Fe,Al)O heterostructures

Chapter 7

132

7.1. Introduction

In recent time, a great interest in the field of spintronics devices widely deals with several

kind of Schottky and heterojunctions of ferromagnetic materials/semiconductors with other non

magnetic semiconductors. Since the rediscovery of colossal magnetoresistance (CMR) effect in

manganite thin films, much attention has been focused on the fabrication of artificially designed

structures [1-3], such as magnetic tunnel junctions, and p–n junctions [4-7] to verify device

concepts based on oxide materials. The doped manganite La1−xAxMnO3 (A=Ca, Sr, and Ba) is a

strongly correlated-electron system with charge, orbital, spin, and lattice degrees of freedom,

possessing diverse physical phenomena. The heterostructures based on p-type perovskite oxides

show some special characteristics such as high magnetic sensitivity, ultraviolet photo voltage,

current field modulations, and the photo carrier injection effect [8-9]. The La0.7Sr0.3MnO3

(LSMO) is a typical double exchange highly spin-polarized system with a high Curie

temperature of about 360 K due to its large one-electron bandwidth. Therefore, it is one of the

best choices as a ferromagnetic electrode. On the other hand, ZnO is an n-type semiconductor

with a wide band gap and large exciton binding energy. However, there have been only a few

reports on the effect of magnetic fields on transport properties in LSMO/ZnO p-n

heterostructures [10,11]. On the other hand ZnO doped with transition metals shows room

temperature ferromagnetic properties, which is also called the dilute magnetic semiconductor

(DMS). With carrier concentration much smaller than the magnetic impurity concentration, the

DMS system provides a complimentary limit to Kondo system. The coupling between localized

impurity spin (S) and mobile valence band holes or conduction band electron can be represented

by the exchange interaction – J S.σ, where σ is the fermion spin operator. The Fe doped ZnO

shows room temperature ferromagnetic behavior and incorporating 1% Al shows enhanced

carrier induced ferromagnetism mainly due to the increase of charge carrier concentrations [12].

A few research work have been carried out with heterojunction of carrier induced ferromagnetic

n-type ZnO DMS with ferromagnetic half metallic p-type LSMO [13]. But a detail study

(thorough investigation) of spin injection through such p-n heterojunction is necessary for

complete understanding to use it in proper spintronics device applications.

In this chapter, we have demonstrated a detail junction magnetoresistive properties of

LSMO/ZnO, LSMO/Zn(Fe)O and LSMO/Zn(Fe,Al)O with different Fe concentrations (5, 7, and

Chapter 7 Electronic‐and magneto transport of La0.7Sr0.3MnO3/ZnO, La0.7Sr0.3MnO3/Zn(Fe)O and La0.7Sr0.3MnO3/Zn(Fe,Al)O heterostructures

133

10%) and established the temperature dependent spin injection and spin extraction process

through the junction.

7.2. Experimental Procedure

7.2.1. Preparation of target

The LSMO powder was synthesized through chemical pyrophoric reaction process where

we have employed stoichiometric mixtures of high purity La2O3 (99.99 %), SrCO3 (99.9+ %) and

Mn(CH3COO)2 (99.0 %). After final grinding and pelletization of LSMO powders, the pelletized

sample was first heated at 800 °C for 12 h, then at 1000 °C for 12 h and at 1200 °C for another

12 h, with intermediate grinding. Final sintering of the LSMO target was carried out at 1200 °C

for 24 h.

The iron doped ZnO powder was first synthesized by solid state reaction process.

Required amount of high purity ZnO, Fe2O3 powder were well mixed with hand grinder

repeatedly and sintered at 450 ºC till the required phase appeared. Required amount of Al2O3

powder were mixed with the ZnO and Fe2O3 powder to dope 1% aluminum in Zinc iron oxide

target. Finally, the Zn(Fe)O and Zn(Fe,Al)O powders were pelletized and sintered at 450 ºC to

use it as the target for pulsed laser deposition (PLD).

7.2.2. Cleaning of substrate

The c-plane (0001) sapphire substrate has been cleaned repeatedly with De-ionized water,

Acetone and Propanol using ultrasonic vibrator. Each cleaning method has been carried out for

20 min.

7.2.3. Preparation of heterojunction

At first the LSMO film was deposited on (0001) well cleaned sapphire substrate at 800 oC and 0.5 mbar O2 pressure. The laser pulse (248 nm KrF laser) of energy density 4 J/cm2 was

applied on the LSMO target for 20 min at a frequency of 10 Hz. The substrate to target distance

was kept at 4 cm in the chamber. After deposition, the film was sintered at the same physical

condition for 1 hr to get well crystalline samples. Then a portion of the substrates were masked

and the ZnO films were grown on the LSMO. The films were grown on LSMO at a substrate

temperature at 450 oC and in ambient oxygen pressures at 10-5 Torr. The excimer laser was used


Chapter 7

134

for 30 min at a laser pulse frequency (repetition rate) of 10 Hz. Electrical contacts were made

with high purity Ag on LSMO film and highly pure In contact with ZnO as ohmic contacts.

7.2.4. Characterization of heterostructure

The structural and surface morphological studies have been carried out using high

resolution x-ray diffraction (Philips pan analytical x-pert), scanning electron microscope (Carl

Zeiss) and atomic force microscope (Nanonics). The electrical properties have been studied out

in details using Keithley 2612 source meter with 1 microvolt resolution and a DMM (Keithley-

2000) along with a temperature controller (Lakeshore-330). The magnetic field was applied in

the current parallel to the plane (CPP) geometry with the magnetic field parallel to the film plane

using a high precision electromagnet (polytronic, model HEM 100).

7.3. Results and Discussion

7.3.1. Structural and surface study

Fig. 7.1.(a) High resolution XRD pattern of LSMO/ZnO heterojunction on sapphire substrate. (b) Zoomed view of HRXRD along with Gaussian fit of (110) plane of LSMO and(201) plane of ZnO(inset). (c) and (d) are the FESEM image of surface and cross sectional view, respectively.

30 31 32 33 34 35

64 66 68 70 72 74 76

Inta

nsity

(a.u

)

2θ(Degree)

2θ

(b)

Inte

nsity

(a.u

)

20 30 40 50 60 70 80

Inte

nsity

(a.u

)

2θ (Degree)

LSM

O (1

0 0

) LSM

O (1

1 0

)Z

nO (1

0 1

) LSM

O (1

1 1

)A

l 2O3 (0

0 6

)Z

nO (1

0 2

)

ZnO

(2 0

1)

(a)

(c)

100 nm

LSMO

ZnO

(d)

200 nm


135

The high resolution x-ray diffraction (HRXRD) pattern, shown in Fig. 7.1(a), has been

carried out using Cu-Kα radiation. Both LSMO film and ZnO films show non-epitaxial behavior

on (0001) c-plane sapphire substrate. The growth of LSMO is in the directions of (100), (110)

and (111) plane and ZnO is in the direction of (101), (102) and (201) planes have been found.

Figure 7.1(b) shows the zoomed view of LSMO (110) and ZnO (201) planes (inset of Fig.

7.1(b)) along with Gaussian fit. It gives the lattice parameters LSMO and ZnO (aLSMO ~ 3.83 Å

and aZnO ~ 3.53 Å). In Fig. 7.1(c) and 7.1(d), the FESEM micrograph of the top view of one of

the LSMO/ZnO heterojunction and the cross sectional FESEM image of the heterojunction have

been shown, respectively. The FESEM micrograph clearly shows the uniformly grown smooth

LSMO and ZnO film with good coverage on sapphire substrate. The average thickness of the

LSMO layer is around ~ 371 nm and ZnO layer is around ~ 124 nm.

Figures 7.2(a) and (b) show the atomic force microscopic (AFM) image of the surface of

LSMO film on (0001) sapphire substrate and 3-d view of the LSMO film, respectively. The

r.m.s. roughness of the LSMO film is found to be ~2 nm. The AFM scan at the junction has been

taken and their 2-d and 3-d image of the junction has been shown in Fig. 7.2(c) and (d),

Fig. 7.2. AFM image of LSMO surface morphology: (a) 2-d and (b) 3-d view. AFM image of LSMO/ZnO heterojunction: (c) 2-d and (d) 3-d view.

0.03 Volts

-0.04 Volts

1.0µm

0.97 Volts

-0.43 Volts

1.20 Volts

-0.47 Volts

1.0µm

(a(b)

(c) (d)


Chapter 7

136

respectively. From the image contrast in Fig. 7.2(d) it is clear that the ZnO layer has been grown

on LSMO film and it confirms the formation of heterojunction.

7.3.2. Electrical properties study

The room temperature current density-voltage (J-V) properties of ZnO/LSMO,

Zn(Fe)O/LSMO and Zn(Fe,Al)/LSMO heterojunctions with 5% Fe have been shown in Fig. 7.3.

Insets (a) and (b) of Fig. 7.3 are the isothermal LSMO/Ag and ZnO/In ohmic contacts,

respectively measured at several different temperatures. These metallic contacts are linear in all

temperatures which confirm that the non linearity of J-V properties originates from ZnO/LSMO

junctions only, not from metal - semiconductor contacts. Doping with 1% Al enhances the carrier

concentration of ZnO films (~1022 /cm3) and causes a very thin depletion layer across the

heterojunctions. So, the reverse saturation current becomes very high and does not show good

rectifying behavior where as ZnO/LSMO, Zn(Fe)O/LSMO shows reasonably good rectifying

behavior as the films have moderate carrier concentrations (~1019-1021/cm3). Large current flows

through the Zn(Fe,Al)O/LSMO junction rather than the ZnO/LSMO or Zn(Fe)O/LSMO

heterojunctions.

Fig. 7.3. Room temperature J-V properties of LSMO/ZnO, LSMO/Zn(Fe)O and LSMO/Zn(Fe,Al)O heterojunction. Upper inset shows the ohmic nature of I-V behavior for LSMO-Ag electrical contact at several temperatures (77-300 K). Lower inset shows the same for ZnO-In electrical contact.

-5 -4 -3 -2 -1 0 1 2 3 4 5

-100

-50

0

50

100

ZnO

Zn(Fe)O

Cur

rent

(mA

/cm

2 )

Voltage(V)

Zn(Fe,Al)O

-3 -2 -1 0 1 2 3-0.4

-0.2

0.0

0.2

0.4 77 K 100 K 150 K 200 K 250 K 300 K

Cur

rent

(mA

)

Voltage (V)

LSMO/AgOhmic contact

(a)

-4 -3 -2 -1 0 1 2 3 4-15

-10

-5

0

5

10

15

(b)

77 K 100 K 150 K 200 K 250 K 300 K

Cur

rent

(mA

)

Voltage (V)

ZnO/InOhmic contact


137

Figure 7.4(a) shows the J-V behavior of Zn(Fe)O/LSMO junction with different Fe

doping percentages ranging from 0 to 10%. Figure 7.4(b) is the J-V behavior of

Zn(Fe,Al)O/LSMO junction with different percentages. The junction J-V characteristics can be

described as [14],

⎟⎟⎠

⎞⎜⎜⎝

⎛ −=

kTJARVe

JJ s

η)(

exp0 (7.1)

where, J0 is the reverse saturation current density. η and Rs are, the ideality factor and junction

series resistance, respectively. A is the junction area. The parameters evaluated from the forward

J-V curves of all those films have been summarized in Table 7.1. The little increase of junction

series resistance with increasing doping concentration reveals the lowering of carrier

concentration in Zn(Fe)O and Zn(Fe,Al)O films.

Figure 7.5(a) shows the temperature dependent J-V characteristics of LSMO/ZnO

heterojunction. Figures 7.5(b), (c) and (d) are the temperature dependent J-V characteristics of

LSMO/Zn(Fe)O heterojunctions with 5, 7, 10% iron doping, respectively. Accordingly, Fig.

7.6(a), (b) and (c) are the junction J-V characteristics of LSMO/Zn(Fe,Al)O with 5, 7 and 10%

Fe. All parameters evaluated from the fitting using Eq. (7.1) have been listed in Table 7.1. The

temperature dependent series resistance for the LSMO/ZnO, LSMO/Zn(Fe)O and

LSMO/Zn(Fe,Al)O junctions with 5% Fe doping have been shown in Fig. 7.6(d). The junction

series resistance is found to decrease with increasing temperature which is generally expected in

p-n junction characteristics.

Fig. 7.4. Room temperature J-V properties of (a) LSMO/ZnO, LSMO/Zn(Fe)O and (b) LSMO/Zn(Fe,Al)O heterojunction with different Fe concentrations.

-8 -6 -4 -2 0 2 4 6-10-505

101520253035

Cur

rent

(mA

/cm

2 )

Voltage (V)

300 KLSMO/Zn(Fe)O

0%5%

7%

10%

H=0 T

(a)-3 -2 -1 0 1 2 3

-100

-50

0

50

100

(b)

Cur

rent

(mA

/cm

2 )

Voltage(V)

5%

10%

7% 300 KLSMO/Zn(Fe,Al)O

H=0 T


Chapter 7

138

Fig.7.5. (a) J-V properties of LSMO/ZnO heterojunction at different isothermal temperatures, (b), (c) and (d) are the same plots for LSMO/Zn(Fe)O heterostructures with 5, 7 and 10% Fe concentrations, respectively.

-10 -5 0 5-10-505

101520253035

77 K100 K

150 K200 K

250 K

Cur

rent

(mA

/cm

2 )

Voltage (V)

300 K

LSMO/ZnO

(a)

H=0 T

-10 -8 -6 -4 -2 0 2 4 6-10-505

1015202530

77 K

100 K150 K

200 K

250 K

Cur

rent

(mA

/cm

2 )

Voltage (V)

300 K

LSMO/Zn(Fe)Owith 5% Fe

(b)

H=0 T

-10 -8 -6 -4 -2 0 2 4 6-10

-5

0

5

10

15

20

25

77 K

100 K

150 K

200 K

250 K

Cur

rent

(mA

/cm

2 )

Voltage (V)

300 K

LSMO/Zn(Fe)Owith 7%

(c)

H=0 T

-10 -8 -6 -4 -2 0 2 4 6

-5

0

5

10

15

20

100 K

150 K

200 K

250 K

Cur

rent

(mA

/cm

2 )

Voltage (V)

300 K

LSMO/Zn(Fe)Owith 10% Fe

(d)

H=0 T

77 K

Fig. 7.6. (a) J-V properties of LSMO/ZnO heterojunction at different isothermal temperatures, (b), (c) and (d) are the same plots for LSMO/Zn(Fe)O heterostructures with 5, 7 and 10% Fe concentrations, respectively.

50 100 150 200 250 300

0.0

0.5

1.0

1.5

2.0

2.5

3.0

ZnO ZnFeO ZnFeAlO

Rs (Ω

-cm

2 )

Temperature (K)

With out field

(d)

-3 -2 -1 0 1 2 3

-100

-50

0

50

100

77 K 100 K 150 K 200 K 250 K 300 K

Cur

rent

(mA

/cm

2 )

Voltage(V)

LSMO/Zn(Fe,Al)O5% Fe

(a)

H=0 T

-3 -2 -1 0 1 2 3-100-80-60-40-20

020406080

100

77 K 100 K 150 K 200 K 250 K 300 K

H=0 T

Cur

rent

(mA

/cm

2 )

Voltage(V)

Zn(Fe,Al)O7% Fe

(b)

-3 -2 -1 0 1 2 3-80-60-40-20

020406080

77 K 100 K 150 K 200 K 250 K 300 K

H=0 T

Cur

rent

(mA

/cm

2 )

Voltage(V)

Zn(Fe,Al)O10% Fe

(c)


139

Fig. 7.7. (a) Isothermal J-V characteristics of LSMO/ZnO heterojunction at several different temperatures under 0.7 T applied magnetic field, (b), (c) and (d) are the same plots for LSMO/Zn(Fe)O heterostructures with 5, 7 and 10% Fe concentrations, respectively.

-9 -6 -3 0 3 6-10-505

101520253035

H=0.7 T

300 K 250 K 200 K 150 K 100 K 77 K

Cur

rent

(mA

/cm

2 )

Voltage(V)

LSMO/ZnO

(a)

-10 -8 -6 -4 -2 0 2 4 6-10

0

10

20

30

H=0.7 T

300 K 250 K 200 K 150 K 100 K 77 K

Cur

rent

(mA

/cm

2 )

Voltage(V)

LSMO/Zn(Fe)O5% Fe

(b)

-10 -8 -6 -4 -2 0 2 4 6-10

-5

0

5

10

15

20

25

H=0.7 T

300 K 250 K 200 K 150 K 100 K 77 K

Cur

rent

(mA

/cm

2 )

Voltage(V)

LSMO/Zn(Fe)O7% Fe

(c)

-10 -8 -6 -4 -2 0 2 4 6

-5

0

5

10

15

20

25 300 K 250 K 200 K 150 K 100 K 77 K

Cur

rent

(mA

/cm

2 )

Voltage(V)

LSMO/Zn(Fe)O10%Fe

(d)

H=0.7 T

50 100 150 200 250 3000.0

0.4

0.8

1.2

1.6

2.0

2.4

2.8

Rsm

(Ω−c

m2 )

Temperature (K)

ZnO ZnFeO ZnFeAlO

With applied magnetic field(d)

H=0.7 T

-3 -2 -1 0 1 2 3-100

-80-60-40-20

020406080

100

H=0.7 T

Cur

rent

(mA

/cm

2 )

Voltage(V)

300K 100K 200K 250K 150K 77K

(a)

LSMO/Zn(Fe,Al)Owith 5% Fe

-3 -2 -1 0 1 2 3

-100

-50

0

50

100

H=0.7 T

300K 100K 200K 250K 150K 77K

Cur

rent

(mA

/cm

2 )

Voltage(V)

Zn(Fe,Al)O7%

(b)

-3 -2 -1 0 1 2 3

-100

-50

0

50

100

H=0.7 T

300K 100K 200K 250K 150K 77K

Cur

rent

(mA

/cm

2 )

Voltage(V)

Zn(Fe,Al)O10%

(c)

Fig. 7.8. (a), (b) and (c) Isothermal J-V characteristics of LSMO/Zn(Fe,Al)O heterostructures at several different temperatures under 0.7 T applied magnetic field with 5, 7 and 10% Fe concentrations, respectively. (b) Temperature dependent series resistance plot for LSMO/ZnO, LSMO/Zn(Fe)O and LSMO/Zn(Fe,Al)O with 5% Fe and 1% Al under 0.7 T applied magnetic field.


Chapter 7

140

Table 7.1. Evaluated fit parameters from the J-V characteristics employing Eq. (1) with and without applied 0.7 T magnetic field.

B

(T)

T

(K)

ZnO/LSMO Zn(Fe)O/LSMO Zn(Fe,Al)O/LSMO

5% 7% 10% 5% 7% 10%

0 n Rs

kΩ

-cm2

J×10-6

mA/cm

2

n Rs

kΩ-

cm2

J ×10-6

mA/cm

2

n Rs

kΩ-

cm2

J× 10-6

mA/c

m2

n Rs

kΩ-

cm2

J× 10-6

mA/cm

2

n Rs

kΩ-

cm2

J×10-6

mA/cm

2

n Rs

kΩ-

cm2

J×10-6

mA/c

m2

n Rs

kΩ-

cm2

J×10-6

mA/cm

2

300 3.1 0.13 7.52 6.9 0.14 5.7 7.2 0.18 9.6 7.2 0.2 8 0.76 0.03 7.5 0.76 0.03 6.8 0.76 0.03 4.6

250 4.1 0.2 3.96 8.2 0.15 3.7 4.1 0.28 8.8 4.1 0.29 7 0.18 0.04 3.7 0.18 0.04 1.7 0.18 0.05 3.19

200 4.0 0.28 3.92 4.0 0.33 1.8 4.0 0.3 7.8 4.0 0.39 5.8 15.5 0.03 3.3 15.5 0.04 3.3 15.5 0.04 2.9

150 22 0.34 3.7 17 0.42 1.6 17 0.46 7.4 17 0.5 4.9 21.4 0.04 2.9 21.4 0.04 7.8 21.4 0.05 2.4

100 10 1.88 3.5 10 2.14 1.2 10 2.32 6.7 10 2.42 4.7 48.3 0.03 2.1 48.3 0.04 6.5 48.3 0.04 1.9

77 106 1.97 7.8 29 2.97 1.1 44 2.88 6.57 59 2.72 2.8 7.4 0.11 2 7.46 0.12 6.1 7.46 0.12 1.8

0.7 300 3.0 0.12 7.7 6.9 0.12 5.8 7.2 0.16 9.7 7.2 0.17 8.5 1.1 0.02 7.8 0.76 0.03 6.2 0.76 0.03 4.7

250 4.1 0.25 3.7 4.1 0.35 2.9 4.1 0.33 7.9 4.1 0.33 6.9 127 0.07 3.5 0.18 0.06 3.6 0.18 0.06 2.9

200 5.1 0.33 3.4 4.0 0.41 1.62 4.0 0.41 7.1 4.0 0.44 5.5 14.9 0.04 3.2 15.5 0.04 8.2 15.5 0.05 2.59

150 22 0.34 3.7 20 0.42 1.61 17 0.45 5.8 17 0.49 4.9 21.4 0.04 3 21.4 0.04 7.08 21.4 0.04 2.58

100 16 1.70 3.1 10 2.06 1.4 10 2.23 6.9 10 2.28 4.5 35.6 0.04 2.6 48.3 0.03 6.08 48.3 0.03 2.3

77 125 1.71 2.7 29 2.7 1.2 44 2.57 5.3 59 2.47 3 22.4 0.17 2.2 7.46 1.02 5.9 7.46 0.09 1.6


141

Figure 7.7(a) shows the temperature dependent J-V characteristics of LSMO/ZnO

heterojunction under 0.7 T applied magnetic field. Figures 7.7(b), (c) and (d) are the temperature

dependent J-V characteristics of LSMO/Zn(Fe)O heterojunctions with 5, 7, 10% iron doping,

respectively under 0.7 T applied magnetic field. Figures 7.8(a), (b) and (c) are the junction J-V

characteristics of LSMO/Zn(Fe,Al)O with 5, 7 and 10% Fe doping under 0.7 T applied magnetic

field. The parameters in Eq. (7.1) are modified under magnetic field (J0, η and Rs are now

replaced by J0m, ηm and Rsm, respectively). The modified value of parameters evaluated

employing modified Eq. (7.1) has also been enlisted in Table 7.1. The temperature dependent

series resistance under 0.7 T magnetic field for the LSMO/ZnO, LSMO/Zn(Fe)O and

Zn(Fe,Al)O junctions with 5% Fe has been shown in Fig. 7.8(d). The junction series resistance

under magnetic field also decreases with increasing temperature.

7.3.3. Junction Magnetoresistance properties

Fig. 7.9. Comparative J-V characteristics of LSMO/ZnO (I), LSMO/Zn(Fe)O (II) and LSMO/Zn(Fe,Al)O (III) heterostructures with and without applied 0.7 T magnetic field. The Fe and Al concentration are 5% and 1%, respectively. (a), (b), (c) and (d) are the same J-V characteristics measured at 77, 200, 250 and 300 K, respectively.

I

II

III

I

II

III

I

II

III

I

II

III

(a) (b) (c) (d)

1 2 3 4 5 60.40.60.81.01.21.41.61.82.0

0T

Cur

rent

(mA

/cm

2 )

Voltage(V)

77 KZnO/LSMO

0.7T

1 2 3 4 5 62468

1012141618

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnO/LSMO at 200K

0.7T

0T

1 2 3 4 5 60

2

4

6

8

10

12

14

16

Voltage(V)

Cur

rent

(mA

/cm

2 )

ZnFeo/LSMO at 200K0T

0.7T

0.0 0.5 1.0 1.5 2.0 2.5 3.00

10

20

30

40

50

60

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnFeAl/LSMO at 200K

0.7T0T

0 1 2 3 4 5 60

5

10

15

20

25

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnO/LSMO at 250K

0.7T

0T

0 1 2 3 4 5 60

5

10

15

20

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnFeO/LSMO at 250K

0.7T

0T

0.0 0.5 1.0 1.5 2.0 2.5 3.00

10

20

30

40

50

60

70

Cur

rent

(mA

)

Voltage(V)

ZnFeAlO/LSMO at 250K

0.7T

0T

0 1 2 3 4 50

5

10

15

20

25

30

35

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnO/LSMO at 300K

0.7T

0T

0 1 2 3 4 50

5

10

15

20

25

30

35

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnFeO/LSMO at 300K

0.7T

0T

0.0 0.5 1.0 1.5 2.0 2.5 3.00

20

40

60

80

100

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnFeAlO/LSMO at 300K

0.7T0T

0 1 2 30

10

20

30

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnFeAl/LSMOat 77K0.7T

0T

0 1 2 3 4 5 60.00.20.40.60.81.01.21.41.61.8

Cur

rent

(mA

/cm

2 )

Voltage(V)

ZnFeO/LSMO 0.7T

0T


Chapter 7

142

The detailed comparative study of J-V with and without applied 0.7 T magnetic field

measured at different temperatures (77, 200, 250 and 300 K) for LSMO/ZnO, LSMO/Zn(Fe)O

and Zn(Fe,Al)O heterojunctions with 5% Fe doping have been shown in Fig. 7.9. Figure 7.9(a I-

III) are the J-V characteristics at 77 K, Fig. 7.9(b I-III) are for 200 K, Fig. 7.9(c I-III) are for 250

K and Fig. 7.9(d I-III) for 300 K. It clearly demonstrates that all the heterojunctions show the

negative junction MR at 77 and 300 K and positive junction MR at 200 and 250 K.

The change of junction MR with applied magnetic fields up to 0.7 T for LSMO/ZnO,

LSMO/Zn(Fe)O, LSMO/Zn(Fe,Al)O heterojunctions with 5% Fe doping at a bias voltage of 3 V

has been shown in Fig. 7.10 (a), (b), (c) and (d) for 77, 200, 250 and 300 K, respectively. The

magnetic field dependent junction magnetoresistance (JMR) behavior in the magnetic field

region up to 0.7 T of LSMO/ZnO, LSMO/Zn(Fe)O and Zn(Fe,Al)O junctions with 5% Fe have

been analyzed using a simple empirical relation as [15],

Fig. 7.10. Plot of junction magnetoresistance with applied magnetic field at (a) 77 K, (b) 200 K, (c) 250 K and (d) 300 K of the LSMO/ZnO, LSMO/Zn(Fe)O and LSMO/Zn(Fe,Al)O heterojunctions with 5% Fe concentration at a bias voltage of 3 V.

(b)


143

βαHJMR = (7.2)

where, α and β are coefficients. The junction magnetoresistance behaviors of other doped

samples have also been investigated employing Eq. (7.2). The obtained best fit values of α and β

for LSMO/ZnO, LSMO/Zn(Fe)O and Zn(Fe,Al)O junctions with 5% Fe are listed in Table 7.2.

The exponent is found to be high (~1) at 200 K and it becomes ~0.5 at higher temperatures for

all these heterojunctions.

Table 7.2. Fit parameters (α and β) evaluated from Eq. (2) for undoped LSMO/ZnO, LSMO/Zn(Fe)O and LSMO/Zn(Fe)O with 5% Fe doping.

Sample Temperature (K) α β

ZnO 77 -7.5 0.27

200 3.1 0.82

250 8.4 0.53

300 -4.5 0.5

Zn(Fe)O 77 -4.7 0.39

200 4.2 0.83

250 13.4 0.53

300 -4.1 0.51

Zn(Fe,Al)O 77 -3.8 0.39

200 5.4 0.83

250 31.1 0.45

300 -0.72 0.51

The origin of junction magnetoresistances in those heterostructures can be best explained by the

spin injection theory of magnetic p-n junction. Spin injection generally occurs in n-side while

spin extraction generally occurs in p-side. Both spin injection and spin extraction becomes large

with applied magnetic field. The current through the magnetic p-n junction depends on magnetic

field because of the spin orbit splitting of magnetic materials [16]. The spin splitting of

conduction bands are created by doping with magnetic ions in non-magnetic semiconductors.

Doping with iron causes a large non equilibrium population of polarized electrons in ZnO

conduction band. This non equilibrium population at magnetic n-side enhances the spin injection


Chapter 7

144

through the space-charge region [17]. It causes a high magnetoresistance according to standard

spin injection theory [18].

The temperature dependent junction MR for LSMO/ZnO, LSMO/Zn(Fe)O and

LSMO/Zn(Fe,Al)O with 5% Fe for three different bias voltages have been shown in Fig. 7.11

(a), (b) and (c), respectively. For three different applied bias voltages the junction

magnetoresistance shows a peak near 250 K. The both spin injection and spin extraction are

sensitive to spin lattice relaxation time which is very much dependent on temperature. The high

value of spin lattice relaxation constant near 250 K [19], causes the enhancement of spin

injection (and spin extraction) through the junction and causes a high positive junction

magnetoresistance. At the higher temperatures the spin lattice relaxation constant drops sharply

and spin injection effect die down through the junctions. The properties of magnetic electrodes

dominates and shows little negative junction magnetoresistance. The change of junction MR

Fig. 7.11. Temperature dependent junction magnetoresistance plot for (a) LSMO/ZnO, (b) LSMO/Zn(Fe)O and (c) LSMO/Zn(Fe,Al)O heterojunctions with 5% Fe concentration at three bias voltages 2.5, 3.5 and 4.5 V. (d) Plot of highest value of the % of junction magnetoresistance (left y-axis scale) at 250 K as well as magnetic moment change in μB/Fe2+ (right y-axis scale) with increasing iron concentration in both LSMO/Zn(Fe)O and LSMO/Zn(Fe,Al)O heterojunctions.


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calculated at 250 K at 2.5 V applied bias voltage for different iron concentrations have been

shown in Fig. 7.11(d). The change of magnetic moment (MS) in Bohr magneton per Fe2+of both

Zn(Fe)O and Zn(Fe,Al)O films with doping percentage have also been shown in the Fig. 7.11(d)

[the dashed lines]. Both the magnetic moment per Fe2+ and junction MR decreases with increase

of Fe concentration. The drop of moment with increasing iron concentrations may be due to the

increasing of antiferromagnetic coupling between Fe pairs which occurs at shorter separation

distances and decrease the spin up population in DMS systems. The decrease of junction MR

with increasing iron concentrations in both Zn(Fe)O and Zn(Fe,Al)O system is most likely due to

the decrease of spin up population caused by higher doping.

7.4. Summary

LSMO/ZnO, LSMO/Zn(Fe)O and LSMO/Zn(Fe,Al)O heterojunctions with 5, 7 and 10%

Fe have been fabricated using pulsed laser deposition technique. The junction J-V properties

have been studied with and without applying magnetic field. The junction with ZnO and

Zn(Fe)O shows good rectifying behavior at all temperatures but heterojunctions with Zn(Fe,Al)O

shows a quick break down at reverse bias. Higher carrier concentration in Zn(Fe,Al)O most

likely causes thin depletion region and hence causes high current transport through the junction.

All the heterojunctions show high positive junction magnetoresistance at a certain temperature

range (150 to 280 K) and low negative magnetoresistance at 77 K and 300 K. The junction

magnetoresistance enhances due to incorporation of 1% Al. This has been best explained using

spin injection theory through magnetic p-n junction. Increase of Fe doping concentration

decreases the junction magnetoresistance as well as magnetic moment per Fe2+ for both Zn(Fe)O

and Zn(Fe,Al)O systems. This is mainly due to the decrease of population of polarized electron

in conduction band. The properties of junction magnetoresistance of such heterostructures can be

extremely useful in the area of spintronic devices.

References

[1] C. Mitra, P. Raychaudhuri, K. Do¨rr, K. H. Mu¨ller, L. Schultz, P.M. Oppeneer, and S.Wirth, Observation of Minority Spin Character of the New Electron Doped Manganite La0.7Ce0.3MnO3 from Tunneling Magnetoresistance, Phys. Rev. Lett. 90, 017202 (2003). [2] C. Mitra, G. Köbernik, K. Dörr, K. H. Müller, L. Schultz, P. Raychaudhuri, R. Pinto, and E. Wieser, Magnetotransport properties of a room temperature rectifying tunnel junction made of electron and hole doped manganites, J. Appl. Phys. 91, 7715 (2002).


Chapter 7

146

[3] C. Mitra, P. Raychaudhuri, G. Köbernik, K. Dörr, K.-H. Müller, L. Schultz and R. Pinto, p-n diode with hole- and electron-doped lanthanum manganites, Appl. Phys. Lett. 79, 2408 (2001). [4] H. Tanaka, J. Zhang, and T. Kawai, Giant Electric Field Modulation of Double Exchange Ferromagnetism at Room Temperature in the Perovskite Manganite/Titanate p-n Junction, Phys. Rev. Lett. 88, 027204 (2002). [5] A. Tiwari, C. Jin, D. Kumar, and J. Narayan, Rectifying electrical characteristics of La0.7Sr0.3MnO3/ZnO heterostructure, Appl. Phys. Lett. 83, 1773 (2003). [6] B. B. Nelson-Cheeseman, F. J. Wong, R. V. Chopdekar , E. Arenholz, and Y. Suzuki, Room temperature magnetic barrier layers in magnetic tunnel junctions, Phys. Rev. B 81,214421 (2010). [7] G. Li , De-bin Huang, Shao-wei Jin , Yong-qing Ma, Xiao-guang Li, Electrical transport properties of heteroepitaxial p_n junction of charge-ordered La7/16Ca9/16MnO3 and 0.5 wt% Nb doped SrTiO3, Solid State Comm. 150, 1737 (2010). [8] H. B. Lu, S. Y. Dai, Z. H. Chen, Y. L. Zhou, B. L. Cheng, K. J. Jin, L. F. Liu, G. Z. Yang and X. L. Ma, High sensitivity of positive magnetoresistance in low magnetic field in perovskites oxide p–n junctions, Appl. Phys. Lett. 86, 032502 (2005). [9] K. Zhao, Y. Huang, Q. Zhou, K. J. Jin, H. Lu, M. He, B. Cheng, Y. Zhou, Z. Chen and G. Yang, Ultraviolet photovoltage characteristics of SrTiO3−δ/Si heterojunction, Appl. Phys. Lett. 86, 221917 (2005). [10] C. Mitra, P. Raychaudhuri, K. Dörr, K. H. Müller, L. Schultz, P. M. Oppeneer and S. Wirth, Observation of Minority Spin Character of the New Electron Doped Manganite La0.7Ce0.3MnO3 from Tunneling Magnetoresistance, Phys. Rev. Lett. 90, 1107202 (2003). [11] C. Mitra, P. Raychaudhuri, G. Köbernik, K. Dörr, K.-H. Müller, L. Schultz and R. Pinto, p–n diode with hole- and electron-doped lanthanum manganites, Appl. Phys. Lett. 79, 2408 (2001) [12] S. Chattopadhyay, T.K. Nath, A.J. Behan, J.R. Neal, D. Score, Q. Feng, A.M. Fox, G.A. Gehring, Temperature dependent carrier induced ferromagnetism in Zn(Fe)O and Zn(FeAl)O thin films, Appl. Surf. Sc. 257, 381 (2010). [13] L. Yan, W. C. Goh, and C. K. Ong, Magnetic and electrical properties of La0.7Sr0.3MnO3– Zn0.8Co0.2Al0.01O junctions on silicon substrates, J. Appl. Phys. 97, 103903 (2005). [14] S. J. May and B. W. Wessels, High field magnetoresistance in p-(In,Mn)As/n-InAs heterojunctions, Appl. Phys. Lett. 88, 072105 (2006). [15] Z. G. Sheng, W. H. Song, Y. P. Sun, J. R. Sun, and B. G. Shen, Crossover from negative to positive magnetoresistance in La0.7Ce0.3MnO3-SrTiO3-Nb heterojunctions, Appl. Phys. Lett. 87, 032501 (2005). [16] J. Fabian, I. Žutić and S. Das Sarma , Theory of spin-polarized bipolar transport in magnetic p-n junctions, Phys. Rev. B 66, 165301 (2002). [17] Mark Johnson, Magnetoelectronics, academic press, Elsevier (2005). [18] S. Chattopadhyay and T. K. Nath, Room temperature enhanced positive magnetoresistance in Pt and carrier induced Zn(Fe)O and Zn(Fe,Al)O dilute magnetic semiconductors junction, J. Appl. Phys. 108, 083904 (2010). [19] A. I. Lobad, R. D. Averitt, C. Kwon, and A. J. Taylor, Spin–lattice interaction in colossal magnetoresistance manganites, Appl. Phys. Lett. 77, 4025 (2000).

Chapter 8

Conclusions

Conclusions Chapter 8

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8.1. Conclusions of thesis

The temperature dependent spin injection or extraction phenomena in magnetic

semiconductors and semimetals have been widely studied in this thesis for possible active

spintronics device applications. The Zn(Fe)O and Zn(Fe,Al)O thin films have been

selected as a room temperature Dilute Magnetic Semiconductor (DMS). It shows n-type

semiconducting behavior with wide band gap. La0.7Sr0.3MnO3 has been chosen as a hole

doped half-metallic manganites as a p-type ferromagnetic electrode. The Zn(Fe)O and

Zn(Fe,Al)O highly crystalline epitaxial thin films show room temperature ferromagnetic

ordering with carrier mediated ferromagnetism. Doping with 1 % Al in ZnFeO film, the

magnetic moment per Fe2+ ion is found to enhance by 3 times as compared to the ZnFeO

film without Al doping. The magnetic moments die down due to higher doping of Fe in

ZnO. The ρ-T behavior shows that the films are semiconducting in nature and the

electronic transport mechanisms in these films have been identified as Variable Range

Hopping in lower temperature range, Efros’s Variable Range Hopping in the intermediate

temperature range and thermally activated transport in higher temperature range.

Analyzing Ordinary Hall Effect data, it is found that the dominant donor has an activation

energy ranging from 33 to 41 meV and several type of scattering mechanisms are present.

Anomalous or Extra-ordinary Hall Effect results show intrinsic ferromagnetic behavior in

these DMS films and both skew scattering and side jump mechanisms are responsible for

the origin of Anomalous Hall voltage in these DMS films. The spin injection from

Zn(Fe)O and Zn(Fe,Al)O Dilute Magnetic Semiconductor to non-magnetic Pt has been

demonstrated explicitly. The spin injection in Pt/Zn(Fe)O and Pt/Zn(Fe,Al)O shows

positive JMR at room temperature and the JMR strictly depends on the magnetic moment

of the films. La0.7Sr0.3MnO3 also shows good ferromagnetic behavior at all temperatures

up to room temperature. The ρ-T behavior of LSMO shows resistivity minima at ~50 K

and a metal-insulator transition peak at ~250 K. The low temperature minima and the

rising of resistivity with the decrease of temperature below 50 K can be best explained

through electron-electron interaction as predicted by quantum interference effect (QIE).

Electron tunneling phenomena through different SiO2 layer in La0.7Sr0.3MnO3/SiO2/p-Si

has also been demonstrated, and, the temperature and oxide defect dependency in positive

JMR is presented. The observed JMR behavior has been explained using tunneling model

Chapter 8 Conclusions

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where the Frenkel-Poole type tunneling mechanism dominates. The temperature

dependent spin injection and spin extraction in La0.7Sr0.3MnO3/ZnO,

La0.7Sr0.3MnO3/Zn(Fe)O and La0.7Sr0.3MnO3/Zn(Fe,Al)O heterostructures have also been

investigated and it shows that the highest spin injection (i.e. maximum JMR) appears at a

temperature ~ 250 K which is the temperature where La0.7Sr0.3MnO3 shows highest spin

relaxation. The junction magnetoresistive properties of these heterostructures have been

best explained using the standard spin injection mechanism through the magnetic p-n

junction. Junction magnetoresistance dies out with the increase of doping concentrations

of Fe in ZnFeO or Zn(Fe,Al)O films for all the three type of heterojunctions due to the

less non equilibrium population of polarized electrons. The positive JMR for these cases

also drastically enhances with enhancing the magnetic moments of Zn(Fe)O films with

incorporating Al.

8.2. Scope of future work

To propose the scope of future work based upon our present dissertation work,

first we would like to consider the study on dilute magnetic semiconductors using other

II-VI and III-V semiconductors. We would like to study the DMS materials using the

magneto-optical methods and would like to search the origin of ferromagnetism in these

dilute magnetic semiconductors.

We would also like to fabricate different kind of homo, hetero and Schottky

junctions using different magnetic metals, semiconductors and half-metals to study the

spin injection properties through them as they have potentials in active spintronic device

applications.

8.3. Contribution of thesis

The temperature dependent spin injection or extraction phenomena involving the

magnetic semiconductors, semimetals, other non-magnetic semiconductor or metals have

been widely investigated for possible active spintronics device applications. As the

semiconductor spintronic devices are still questionable in the area of both basic science

and technology, this thesis can lead to various possible ideas to design different

spintronics heterojunctions for industrial applications.

Spin Tronic s

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