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Preparation of Giant Magnetic Materials of

Iron Nitride Related Compounds

Ph.D. Thesis

Shahid Atiq

Roll No. P0403

Session (2004-2009)

CENTRE OF EXCELLENCE IN SOLID STATE PHYSICS

UNIVERSITY OF THE PUNJAB

LAHORE (PAKISTAN)

Preparation of Giant Magnetic Materials of

Iron Nitride Related Compounds

A thesis submitted, in partial fulfillment

of the requirement for the award of the degree of

DOCTOR OF PHILOSOPHY

In

SOLID STATE PHYSICS

By

Shahid Atiq

Session (2004-2009)

Centre of Excellence in Solid State Physics

University of the Punjab

Lahore, Pakistan

“Dua”

CERTIFICATE

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

Atiq S/o Muhammad Rafique, Roll No. P-0403, Session (2004-2009), as partial

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

this thesis to Centre of Excellence in Solid State Physics, University of the Punjab, Lahore.

Research Supervisor Director

Dr. Saadat. Anwar Siddiqi, Dr. Saadat Anwar Siddiqi

Professor, Professor and Director,

Centre of Excellence in Centre of Excellence in

Solid State Physics Solid State Physics

University of the Punjab, University of the Punjab,

Lahore. Lahore.

Dedication

Success always solicits

for two things

Exertion & Fortune

If I am successful then my exertion is the efforts of my Parents

which they made to fulfill my wishes and my fortune is due to

their prayers. Hence I dedicate my success to my

Beloved Parents

Acknowledgements

Allah Almighty has blessed me with countless blessings. Without His Will, nothing

is possible. I glorify Allah Almighty and send salutations upon His last Prophet

Muhammad (PBUH), and praise and venerate Allah Almighty and send to him. All praise

is due to Allah Almighty, whose favors are sole cause for enabling me to accomplish this

tedious research work.

I would like to offer the most sincere thanks to my highly respected, the honorable,

and kind research supervisor Prof. Dr. Saadat Anwar Siddiqi for his all time guidance,

tremendous co-operation and above all, the kind hearted sympathetic attitude which he

rendered towards me from the day I took admission in Ph.D. Dear Sir! Hats off to you.

The higher education commission of Pakistan is highly acknowledged for

supporting this research work through Indeginous-5000 Fellowship Program and support

through International Research Support Initiative Program (IRSIP).

I express my deepest sense of gratitude to Prof. Dr. Nazma Ikram, for her nice

parental treatment, valuable advices and timely help, whenever needed. Thanks must go to

Prof. Dr. Shahzad Naseem for his co-operation as I have disturbed him so many times for

technical helps, and every time I found him ready to help. An ideal for everyone, Prof. Dr.

Tariq Abdullah, from whom I have learnt a lot, not only academic knowledge but also the

ways and manners of a disciplined life. I pray for his long and peaceful life. I am also

pleased to offer sincere thanks to my teachers, Prof. Dr. F. M. Nazar and Dr. Saira Riaz.

I don‟t want to be late to acknowledge some special one‟s who really changed my

life, so let‟s move to KAIST, South Korea. The limitless help, encouragement, and

guidance of Prof. Dr. Sung-Chul Shin made it possible to do a quality work by contributing

his wide research experience. He is a genius person, who took care of everything of mine,

including even meals, housing etc. It‟s my privilege to express my gratitude to him. My

sincere gratitude also goes to Dr. Tayyab Imran, from whom everything started, who

guided us to KAIST and arranged recommendations for us. Dr. Tayyab, thank you!

The days I spent at KAIST with a highly professional group of researchers like Dr.

Viswan Ravindranath, Dr. Jae-Woo Jeong, Dr. Kwang-Su Ryu, Dr. Kyeong-Dong Lee, Dr.

Hyun-Suk Ko, Young-Sung Park, Sang-Hyun Kim, Ji-Wan Kim, Sung-Hyun Lee, Chang-

Yup Park, and Hyun-Sung Lee are really an asset of my life. I am grateful to all of them as

they never made me feel alone at abroad. I was deeply impressed by the sharp „presence of

mind‟ of young Dr. Hyun-Suk Ko during work on Magnetron Sputtering. Special thanks

are due to Dr. Ravi for his technical tips, Sang-Hyun for helping on AFM and MFM

results, Hyun-Sung for helping on XRDs, and Sung-Hyun for taking care of everything

else. I would mention here, as experienced at KAIST, the weekends are the best days to do

research work.

Now, the genius of Pakistan, a role model for all of us, Dr. Sabeih Anwar…what a

personality! He is a source of inspiration and motivation for me, and this is what we need

to become a talented researcher. I can just say, may God bless you Dr. Sabeih!

My humblest thanks must go to Prof. Dr. Beong-Ki Cho for providing me

internship opportunity at GIST, South Korea. Some memorable days of my life I have

spent at GIST in the nice company of Young-Man, Ju-Young, Seung-Ha, Ki-Su, Moon-

Jung, Tallal, Irfan, Hamayun, Zafar, Abdul-Baqi, Furrukh, Naeem, Tauseef and Mudassar.

Thanks must go to Shahid Mahmood Ramay for his time to time help, both at home

and abroad. I also acknowledge the nice company of Furrukh Shahzad, Huma Pervaiz,

Muhammad Zaffar, Qamar-u-Zaman, Fazal Abbas, Adeel Riaz, Uzma Ghazanvi, Murtaza

Saleem, and Muhammad Tanvir. Thank you all for your moral support. Dear Zaffar, your

timely help at many occasions is unforgettable for me.

Qamar Zia, Rafi Raza, and I became friends on first sight. Some unforgettable

moments I have spent with them. Dear friends! Your help was of special kind and so my

thanks are. I have warm feelings of love and affection for Asim, Farooq, Kamran, Junaid,

Aman Ullah, Sami Ullah, Yasir Saeed, Saeed , and Gul Bahar for their moral support. Dear

Shahbaz Hayat, very warm thanks are offered to you too, for being critical and correct.

Once more I have to switch over to KAIST to recognize the support of a distinct

group, the Kaistian-Pakistanis, for their continuous help during my stay there. Dear Fiaz!

Thank you for your help in understanding the Korean way of life. Dear Dr. Ehsan,

Dr. Raza, Dr. Faisal, Dr. Maqsood, Dr. Zohair, Aamir Bhai, Anwar Bhai, Asif, and Nabeel

your hospitality was outstanding that I‟ll never forget.

Thank you my dear brothers, Zahid Rafique, Sajid Rafique, Majid Rafique, Fahad

Rafique and dear sisters for your endless encouragement, co-operation and prayers for my

success. Thanks must go to my cousins Shaukat Ali, Muhammad Javed, Muhammad

Ramazan, Abbas Ali and all others for their continuous support. Brother Javed, thank you

for encouraging me from the day I went to school. I pray for a long healthy life for my dear

uncle Muhammad Ishaque, who always cared about my future more than anyone else.

I can never forget my school teachers, Muhammad Anwar Kazmi, Yousaf Masih

Gull, and M. Aslam Khokhar who really shaped my educational carrier from the very

basics. The humblest thanks are offered in respect of my school teachers. I am deeply

grateful to Ch. Noor Muhammad, Hamid Saeed Bhatti, M. Akmal, M. Asghar, Javed

Aslam, and Liaqat Ali Gill for their all time guidance in my professional carrier.

True friends are a real wealth. I thank all of my friends and feel proud to have a nice

company of sincere friends like Nawazish Ali, Zahid Baig, Muhammad Anwar, Ehsan-ul-

Haque, Muhammad Shahbaz, Shahbaz Zaki, Tariq Najam, Asim Rafique, and Sh. Imran,

who always pray for me. Very special thanks to Muhammad Anwar, who kept me in his

prayers even in Bait Ullah and Masjad-Nabvi.

This acknowledgement cannot complete without mentioning some unseen persons,

they are my pupils and their parents who are always praying for me. Thanks to you all. I

feel lucky for that. I am also thankful to all the staff members of library, laboratory, and

clerical office.

How can I forget my parents, who played a vital role in my grooming. I extend my

heartiest complements to my parents whom I am dedicating this thesis. My fervent thanks

are to my mother. She is bestowed with exemplary motherhood manners. I know my wife

and kids, Kinza, Haroon, and Hassan have missed me a lot during my research period.

Their sacrifice can‟t be acknowledged by mere words.

In the end, thanks to all who deserve but remained unmentioned. Thanks to all those

who even think about me.

SHAHID ATIQ

Abstract

Iron nitride compounds have been studied extensively during the last few decades

due to their remarkable magnetic properties which make them a strong candidate for

applications in high density magnetic recording heads and recording media. The magnetic

properties of iron nitride thin films mainly depend on the phase composition which is rather

complex and has a variety of meta stable phases. Two ferromagnetic phases of the Fe-N

system are of special interest. The first, γ′-Fe4N has a cubic structure and contains 20 at%

nitrogen. The other, α″-Fe16N2 can be regarded as an ordered solution of N (about 11 at%)

atoms in tetragonally distorted α-Fe lattice resulting in a bct structure.

We have deposited γ′-Fe4N thin films on different single crystal substrates under

various deposition conditions using magnetron sputtering. Pure target of α-Fe was dc

sputtered, while pure Ar and N2 were used as working and reactive gases respectively,

under a base pressure better than 2 10-6

Torr.

X-ray diffraction (XRD), pole figure and phi scan revealed highly ordered single

phase epitaxial growth of the γ′-Fe4N films deposited on MgO(100) substrates. The

crystalline quality was directly influenced by substrate temperature and annealing time.

The rocking curve analysis verified the highly oriented texture of the films. It was found

that a 450 ˚C substrate temperature and in-situ annealing of 30 minutes were the most

favorable conditions for epitaxial growth of the films.

The square nature of magnetic hysteresis loops manifested strong ferromagnetic

behavior of the films. Angle dependent hysteresis loops showed <100> and <110> as easy

and hard axes of magnetization, respectively. The values of full width at half maximum

(FWHM) and root mean square (RMS) values of surface roughness were found to affect

saturation magnetization (Ms) very strongly. The lowest value of surface roughness (0.17

nm) and minimum FWHM value of 0.37° were characteristic of high quality epitaxial

growth. The Ms was enhanced by both, the decrease in the FWHM and the RMS value of

surface roughness. A further increase in Ms was noticed when the temperature was

decreased to 10 K. The strong ferromagnetic behavior was also witnessed from the static

and dynamic magnetic domain images of γ′-Fe4N films. A quite dense mictrostructure and

the presence of cyclic holographic vertices depicted strong exchange coupling among

magnetic domains. The films exhibited metallic behavior as the resistivity decreased by

decreasing temperature below room temperature.

To investigate the effect of lattice mismatch on the Ms, the γ′-Fe4N films were

deposited on three different single crystal substrates having lattice mismatches from 11% to

0%. The maximum Ms value 1980±20 emu/cc (24% higher than reported for this phase)

was observed for the film deposited on LaAlO3(100) substrate with 0% lattice mismatch

which was a direct evidence that minimum lattice mismatch had guaranteed epitaxy which

in turn enhanced Ms.

α″-Fe16N2 thin films were deposited on Si(100) substrates with varying nitrogen

partial pressure. The films were deposited at a substrate temperature of 200 ˚C and in-situ

annealed for 1 hour. As revealed by XRD, the epitaxial growth of α″-Fe16N2 films was

achieved at a nitrogen partial pressure of 0.8 m Torr, while changing the pressure below

and above this value resulted in the multi-phase iron nitride structure. In this series, the

highest value of Ms was also achieved for the epitaxially grown α″-Fe16N2 sample.

In a search of more stabilized α″-phase, the influence of Co, Pt, and Cr alloying

addition on the saturation magnetization of α″-Fe16N2 thin films has also been investigated.

The at% concentration of these elements in α″-phase was controlled by placing small chips

on iron target during deposition. For Co addition, the Ms increased and reached a maximum

value of 1660±20 emu/cc as the Co concentration reached to a value of 12.86 at%, while

the Ms decreased with further increase in Co at%. When Pt was added to the α″-phase, the

maximum Ms was achieved corresponding to 27.46 at% of Pt. Cr could not favor its

addition, as the Ms decreased straightforwardly, as the Cr contents increased in iron nitride

matrix.

The magnetization in a ferromagnetic system reverses with a sequence of discrete

and jerky jumps, known as the Barkhausen avalanche. Probing into this aspect of

ferromagnetic films, the present study has also focused on the Barkhausen avalanche for its

critical scaling behavior revealing a power-law distribution of the Barkhausen jump size.

The most interesting unsolved question in this field is whether the universality in the

critical exponent holds regardless of the number of defects in ferromagnetic thin films of

γ′-Fe4N. To answer this question, we have investigated the power-law scaling behavior of

the Barkhausen avalanches in epitaxial and non-epitaxial thin films of iron nitride using the

Kerr microscope, capable of time-resolved domain observations. The critical exponent in

the Barkhausen avalanche observed for the epitaxial and non-epitaxial films exhibited

universality, irrespective of the fact that non-epitaxial films were supposed to have a large

number of defects as compared to epitaxial films.

CONTENTS

Chapter 1 MAGNETISM AND THIN FILMS

1.1 Ferromagnetism 1

1.2 Basic Magnetism Concepts 2

1.2.1 Magnetic Susceptibility 2

1.2.2 Magnetic Permeability 4

1.2.3 Magnetic Coercivity 4

1.2.4 Magnetic Anisotropy 5

1.3 Soft Magnetic Materials 7

1.4 Surveying High Magnetic Moment Soft Materials 9

1.4.1 Iron 9

1.4.2 Fe-Si Alloys 9

1.4.3 Fe-Ni Alloys 9

1.4.4 Fe-Al and Fe-Al-Si Alloys 10

1.4.5 Soft Ferrites 10

1.4.6 Amorphous Soft Magnetic Alloys 11

1.4.7 Nano-Crystalline Alloys 11

1.5 Iron Nitride (FeN) Based High Density Head Materials 11

1.5.1 FeN-Based Amorphous Films 12

1.5.2 FeN-Based Nanocrystalline Films 12

1.5.3 FeN-Based Nitride Multilayers 12

1.6 Thin Film Processing 13

1.6.1 Growth Process 13

1.6.2 Factors Affecting Film Properties 13

1.7 Sputtered Films and Recording Applications 14

1.8 Brief Overview of This Dissertation 15

Chapter 2 IRON NITRIDE AND RELATED COMPOUNDS

2.1 Iron Nitride 17

2.2 γ′-Fe4N 18

2.3 α″-Fe16N2 24

2.4 α″-(Fe,X)16N2 28

2.5 Polycrystalline/Multiphase Iron Nitrides 29

Chapter 3 EXPERIMENTAL TECHNIQUES

3.1 Magnetron Sputtering 31

3.1.1 The Sputter Principle 31

3.1.2 The Sputter Parameters 32

3.1.3 The Sputter System 32

3.2 Sample Preparation 34

3.2.1 Cleaning of Substrates 34

3.2.2 Preparation of γ′-Fe4N Thin Films 36

3.2.3 Preparation of α″-Fe16N2 Thin Films 37

3.3 Crystal Phase Determination 37

3.3.1 X-ray Diffraction 37

3.3.2 Rocking Curve 39

3.3.3 Pole Figure and Phi Scan 39

3.4 Atomic Force Microscope 41

3.5 Scanning Electron Microscope 43

3.6 Magnetic Force Microscope 43

3.7 Vibrating Sample Magnetometer 45

3.8 Superconducting Quantum Interference Device 49

3.9 Auger Electron Microscope 50

3.10 Transmission Electron Microscope 52

3.10.1 TEM-Microstructure 52

3.10.2 Lorentz Microscopy 54

3.10.3 Electron Holography 54

3.11 Magneto-Optic Microscope Magnetometer 56

Chapter 4 PREPARATION AND CHARACTERIZATION OF

γ′-Fe4N THIN FILMS

4.1 Preparation of γ′-Fe4N Thin Films 60

4.2 Crystal Phase Determinations 64

4.2.1 X-ray Diffraction Analysis 64

4.2.2 Rocking Curve Analysis 70

4.2.3 Phi Scan Analysis 74

4.2.4 Pole Figure Analysis 74

4.3 Surface Morphology 77

4.4 RMS Surface Roughness 83

4.5 Magnetic Characterization 89

4.5.1 Saturation Magnetization 89

4.5.2 FWHM Dependent Ms 89

4.6 Substrate Dependent Preparation of γ΄-Fe4N 93

4.6.1 Preparation 93

4.6.2 XRD Analysis 93

4.6.3 Substrate Dependent Ms 96

4.7 Magnetic Anisotropy 100

4.8 Static Magnetic Domains 103

4.9 TEM-Microstructure 105

4.10 Lorentz Microscopy and Electron Holography 105

4.11 Low Temperature Magnetic Behavior 109

4.12 Electrical Behavior of γ′-Fe4N Films 109

Chapter 5 PREPARATION AND CHARACTERIZATION OF

α″-Fe16N2 THIN FILMS

5.1 Preparation of α′′-Fe16N2 Thin Films 117

5.2 Crystal Structure Determinations 120

5.3 Structural Morphology of Surface 122

5.4 RMS Roughness of Surface 122

5.5 Magnetic Domain Structure 126

5.6 Saturation Magnetization 126

5.7 Motivation for Studying the α′′-(Fe,X)16N2 System 133

5.8 FeCoN Alloy System 134


5.8.2 Co Composition Determination 134

5.8.3 Effect of Co Addition on Ms 135

5.9 FePtN Alloy System 140


5.9.2 Pt Composition Determination 140

5.9.3 Effect of Pt Addition on Ms 140

5.10 FeCrN Alloy System 146


5.10.2 Cr Composition Determination 146

5.10.3 Effect of Cr Addition on Ms 146

Chapter 6 POWER LAW SCALING BEHAVIOR IN 2D

FERROMAGNETIC THIN FILMS

6.1 Magnetic Domains 151

6.2 Domain Wall Motion 152

6.3 Magnetization Reversal 152

6.3.1 Magnetization Reversal by Spin Injection 155

6.3.2 All-optical Switching 155

6.3.3 Magnetization Reversal in an Applied Field 155

6.4 The Barkhausen Effect 156

6.5 The Critical Scaling Behavior 156

6.6 The Critical Scaling Behavior in 2D Ferromagnetic Films 159

6.7 The Critical Scaling Behavior in 2D γ′-Fe4N Thin Films 159

Chapter 7 CONCLUSIONS 169

REFERENCES 172

LIST OF FIGURES

Figures Contents Page

1.1 Typical M-H loops of soft and hard magnetic materials 3

1.2 (a) A typical magnetic hysteresis loop revealing basic magnetic quantities 6

1.2 (b) M – H loop of bcc iron with easy and hard axis of magnetization 6

1.3 Free pool distributions in magnetized bodies depending on shape anisotry 8

2.1 Crystal Structure of γ′-Fe4N 20

2.2 Schematic changes in the electronic structure at first and second

neighbor iron site, induced by hybridization with nitrogen 21

2.3 Crystal Structure of α″-Fe16N2 25

3.1 Schematic diagram of a dc magnetron sputtering system 33

3.2 Basic features of a typical x-ray diffraction experiment 38

3.3 Definitions of angles in a rocking curve measurement set up 40

3.4 Schematic drawing of an atomic force microscope 42

3.5 Electron optics of a scanning electron microscope 44

3.6 (a) Working mode of a schematic MFM with, (b) close image of

magnetic tip and sample surface 46

3.7 Schematic of a typical vibrating sample magnetometer 48

3.8 Schematic of auger electron spectroscope 51

3.9 Typical components of a Transmission Electron Microscope,

equipped with a Lorentz lens 53

3.10 Schematic ray diagram illustrating experimental geometry used for

electron holography in TEM 55

3.11 Schematic of a Magneto-Optical Microscope magnetometer 57

4.1 XRD patterns of (a) Pure MgO(100) substrate, (b) Pure α-Fe

deposited on MgO(100), Iron Nitride samples deposited using

nitrogen partial pressure of (c) 0.2 mT, (d) 0.4 mT, (e) 0.6 mT,

(f) 0.8 mT, and (g) 1.0 mT 66

4.2 XRD patterns of (a) Pure MgO(100) substrate, γ΄-Fe4N thin films

deposited on MgO(100) substrates at deposition temperatures of

(b) 200 ˚C, (c) 250 ˚C, (d) 300 ˚C, (e) 350 ˚C, (f) 400 ˚C, (g) 450 ˚C,

and (h) 500 ˚C 67

4.3 XRD patterns of γ΄-Fe4N thin films deposited on MgO(100)

substrates at deposited at 450 ˚C and in-situ annealed (a) 10 min.

(b) 20 min. (c) 30 min. (d) 40 min. and (e) XRD pattern of pure

MgO(100) substrate 69

4.4 Rocking curve of γ′-Fe4N films obtained for (100) reflection for the

sample annealed for 10 minutes 71







4.8 FWHM plotted against annealing time of γ′-Fe4N thin films 73

4.9 The phi (ϕ) scan for (200) reflection of γ′-Fe4N films obtained for

the sample annealed for 30 minutes 75

4.10 The pole figure for (200) reflection of γ′-Fe4N films obtained for the


4.11 Scanning Electron Microscope image of γ′-Fe4N sample deposited

at 200 ˚C substrate temperature 78









4.16 Cross-sectional SEM image of γ′-Fe4N sample in-situ annealed for

10 min 81


20 min 81


30 min 82


40 min 82

4.20 Topographical image obtained using AFM for the sample annealed

for 10 min 85


for 20 min 85


for 30 min 86


for 40 min 86

4.24 RMS roughness plotted against annealing time of γ′-Fe4N thin films 87

4.25 RMS roughness and FWHM plotted against annealing time

of γ′-Fe4N thin films 88

4.26 M-H loops obtained using VSM for the samples annealed for

(a) 10 min., (b) 20 min., (c) 30 min., and (d) 40 min 90

4.27 Effect of annealing time plotted against RMS roughness and saturation

magnetization of the γ΄-Fe4N films 91

4.28 XRD patterns of (a) MgO(100) substrate, (b) γ΄-Fe4N on MgO,

(c) SrTiO3(100) substrate, (d) γ΄-Fe4N on SrTiO3, (e) LaAlO3(100)

substrate, and (f) γ΄-Fe4N on LaAlO3 95

4.29 M-H loops for the γ′-Fe4N thin films deposited on (a) MgO(100)

substrate, (b) SrTiO3(100) substrate, and (c) LaAlO3(100) substrate 97

4.30 (a) Ms plotted against FWHM for the samples annealed for 10 – 40

minutes, (b) Ms plotted against % lattice mismatch for the samples

deposited on three different substrates 98

4.31 Hysteresis loops obtained using MOMM for the sample annealed for

(a) 10 min. (b) 20 min. (c) 30 min. and (d) 40 min 101

4.32 Angle dependent hysteresis loops obtained using MOMM for the


4.33 Magnetic domain images captured using MFM for the γ′-Fe4N film

deposited at 450 ˚C and in-situ annealed for 30 minutes 104

4.34 Micro-structural images obtained using TEM for the sample deposited

at 450 ˚C and in-situ annealed for 30 minutes 106

4.35 Micro-structural images obtained using TEM for the pure iron deposited

on MgO(100) 106

4.36 Micro-structural images obtained using TEM for the sample deposited at

450 ˚C and in-situ annealed for 30 minutes in (a) TEM mode, and (b)

Lorentz mode 107

4.37 Holographic images obtained using TEM for the sample deposited at

450 ˚C and in-situ annealed for 30 minutes in (a) high magnification

and (b) low magnification 108

4.38 M-H loops obtained at 300 K, 250 K, 200 K, and 150 K for the sample

deposited at 450 ˚C and in-situ annealed for 30 minutes 110

4.39 M-H loops obtained at 100 K, 50 K, and 10 K for the sample deposited at

450 ˚C and in-situ annealed for 30 minutes 111

4.40 A graph plotted for Temperature against Coercivity for the γ′-Fe4N

thin films 112

4.41 Normalized magnetization plotted against temperature for the

sample deposited at 450 ˚C and in-situ annealed for 30 minutes 113

4.42 Coercivity plotted against Temperature for the the sample deposited at

450 ˚C and in-situ annealed for 30 minutes 115

5.1 XRD patterns of (a) pure Si(100) substrate, iron nitride samples

deposited on Si(100) using nitrogen partial pressure of (b) 0.2 mTorr,

(c) 0.4 mTorr, (d) 0.6 mTorr, (e) 0.8 mTorr, and (f) 1.0 mTorr 121

5.2 Microstructural image obtained by SEM for the iron niride sample

deposited using2NP = 0.4 m Torr 123


deposited using2NP = 0.6 m Torr 123


deposited using 2NP = 0.8 m Torr 124


deposited using 2NP = 1.0 m Torr 124

5.6 Effect of nitrogen partial pressure on the RMS value of surface

roughness 125

5.7 Magnetic domain structure obtained using MFM for iron nitride

sample deposited using nitrogen partial pressure of 0.8 mTorr 127

5.8 Surface topography for iron nitride sample deposited using nitrogen

partial pressure of 0.8 mTorr 127

5.9 M – H loops obtained using VSM for the iron nitride samples deposited

using nitrogen partial pressure of (a) 0.2 m Torr, (b) 0.4 m Torr,

(c) 0.6 m Torr, (d) 0.8 m Torr, and (e) 1.0 m Torr 129

5.10 Effect of nitrogen partial pressure on saturation magnetization 130

5.11 Effect of nitrogen partial pressure on saturation magnetization and

RMS value of surface roughness for iron nitride thin films 132

5.12 AES scans obtained for FeCoN films deposited by varying Co chips 136

5.13 M – H loops obtained by VSM for iron nitride films deposited using

(a) No Co chip, (b) 1 Co chip, (c) 2 Co chips, (d) 3 Co chips, and

(e) 4 Co chips 138

5.14 Effect of no. of Co chips on at% Co in α″-Fe16N2 and

on saturation magnetization 139

5.15 AES scans obtained for FePtN films deposited by varying Pt chips 142


(a) No Pt chip, (b) 1 Pt chip, (c) 2 Pt chips, (d) 3 Pt chips, and

(e) 4 Pt chips 144

5.17 Effect of no. of Pt chips on at% Pt in α″-Fe16N2 and


5.18 AES scans obtained for FeCrN films deposited by varying Cr chips 147


(a) No Cr chip, (b) 1 Cr chip, (c) 2 Cr chips, (d) 3 Cr chips,

and (e) 4 Cr chips 149

5.20 Effect of no. of Cr chips on at% Cr in α″-Fe16N2 and


6.1 (a) (a) Domain formation in a saturated magnetic material driven by

the magnetostatic (MS) energy in the single domain state,

introduction of 180o domain walls reduces the MS energy but

raises the wall energy creating (b) double domain state, (c) multi-

domain state, and (d) 90o closure domains eliminate MS energy but

increase anisotropy energy in a uniaxial material 153

6.1 (b) The internal structure of 180˚ domains 153

6.2 The rotation of the magnetization vector in the (a) Bloch wall,

and (b) Neel wall 154

6.3 Series of local displacement having discrete and jerky jumps during

magnetization reversal revealing Barkhausen effect 157

6.4 The power law scaling behavior 158

6.5 A Series of six domain images showing Barkhausen avalanche in

the film annealed for 10 min 161







6.9 The distribution of Barkhausen avalanche size for the sample

annealed for 10 min 164







6.13 Effect of annealing time on critical exponent τ

observed in 2D iron nitride thin films 168

LIST OF TABLES

Tables Contents Page

2.1 Iron sites in γ′-Fe4N and their nearest neighbors with distances 20

4.1 Nitrogen partial pressure dependent preparation of γ′-Fe4N films 61

4.2 Substrate temperature dependent preparation of γ′-Fe4N films 62

4.3 Annealing time dependent preparation of γ′-Fe4N films 63

4.4 A table showing annealing time and the corresponding values of

RMS roughness, FWHM, Ms, and μB/Fe 92

4.5 Substrate dependent preparation of γ′-Fe4N films 94

4.6 A table showing % lattice mismatches between substrates and the

deposit along with the Ms and μB/Fe values for the γ′-Fe4N films 99

5.1 Nitrogen partial pressure dependent preparation of α″-Fe16N2 119

5.2 The Influence of nitrogen partial pressure on phase formation and

the corresponding Ms value 131

5.3 The elemental composition of Co in iron nitride corresponding to

no. of Co chips 137

5.4 The elemental composition of Pt in iron nitride corresponding to

no. of Pt chips 143

5.5 The elemental composition of Cr in iron nitride corresponding to

no. of Cr chips 148

Chapter 1 Magnetism and Thin Films

1

Magnetism and Thin Films

Thanks to its fascinating properties in both macroscopic and atomic dimensions,

magnetism has attracted the attentions of philosophers and scientists from ancient times

to the present day. Since the first application of magnetite as a compass in pre-historic

China, and from the early middle ages in Europe, magnetic materials have become an

indispensible part of our daily life. Among the most significant applications of

magnetism in the modern world, it is no surprise that these applications deal with

ferromagnetism. Ferromagnetism is easily the most important technological branch of

magnetism and helps further our understanding of the phenomenon of magnetism. Thus,

we can fabricate new magnetic materials with improved properties and make better use of

existing materials as well. This chapter mainly focuses on the basic underlying concepts

of ferromagnetism, ferromagnetic materials and their potential applications in recording

media. The knowledge of thin film processing is essential in order to understand

magnetic thin films, which is described in latter part of this chapter.

1.1 Ferromagnetism

Ferromagnetism is a phenomenon exhibited by certain materials like iron, nickel

or cobalt that become magnetized when placed in a magnetic field. These materials

contain unpaired electrons, each with a small magnetic field of its own, capable to align

readily with one another in response to an external field. This alignment tends to persist

even after the field is removed, the phenomenon is known as magnetic hysteresis.

Broadly speaking, ferromagnetic materials are characterized by a long-range

ordering of their atomic moments, even in the absence of any external field. The

spontaneous, long-range magnetization of a ferromagnetic material is observed to vanish


2

above an ordering temperature called the Curie temperature (Tc). Ferromagnets are

useful because a large magnetic induction (B ≈ 1-2 T) is produced by fairly a small

applied field (H ≈ 100 A/m).

In a ferromagnetic material, there are certain regions called magnetic domains

(ranging in size upwards from approximately 0.1 μm), over which all magnetic moments

(μm) are essentially parallel. Magnetic domains are separated from each other by domain

walls, surfaces over which the orientation of μm changes relatively abruptly (within about

10-100 nm). The magnetizations in different domains have different directions so that

their vector sum may vanish, rendering the overall bulk sample unmagnetized.

1.2 Basic Magnetism Concepts

1.2.1 Magnetic Susceptibility

The magnetic properties of a material are characterized not only by the magnitude

and sign of Magnetization (M) but also by the way in which M varies with the applied

field (H).

The ratio of these two quantities is called the susceptibility (κ):

H

M emu/cm

3 Oe

Since M is the magnetic moment per unit volume, it is sometimes called the

volume susceptibility. In any case, it is well named, since it indicates „how responsive a

material is to an applied magnetic field‟. At large values of H, the magnetization M

becomes constant and reaches a value, called saturation magnetization (Ms). After

saturation, a decrease in H to zero does not reduce M to zero, this phenomenon is named

as hysteresis [1].

If a small applied field suffices to produce saturation, the material is said to be

magnetically soft. On the other hand, magnetically hard materials require a large applied

field to attain the saturation point. The typical hysteresis loops of soft and hard materials

are shown in Figure 1.1. Note that, cgs units of magnetism are used in this section, to

describe these basic magnetic phenomena.


3

Figure 1.1 Typical M-H loops of soft and hard magnetic materials


4

1.2.2 Magnetic Permeability

A physicist is mainly interested in the variation of M with H, because it

immediately gives the susceptibility and discloses the kind of substance. The engineer, on

the other hand, concerned with ferro- or ferrimagnetic materials, wants to know the total

flux density B produced by a given field, known as a B-H curve, also called a

magnetization curve. The ratio of these two quantities is the permeability (μ):

𝜇 = 𝐵 𝐻 … … … … … … … … … … … … (1.1) Since B = H + 4πM, so we have

𝐵 𝐻 = 1 + 4𝜋 𝑀 𝐻 … … … … … … … … … … … … (1.2)

=> 𝜇 = 1 + 4𝜋𝜅… … … … … … … … … (1.3)

„μ‟ is the slope of the B-H curve at a particular value of H. For ferro- and ferrimagnetic

materials both 𝜅 and 𝜇 are large and positive, and both are functions of H. The initial

induction B produced in response to a small field H defines the initial permeability, μi =

(𝐵 𝐻) H≈0. At higher fields B increases sharply and the permeability increases to its

maximum value, μmax.

The induction and magnetization that remain in the sample when the applied field

is zero are called the residual induction (Br) and remanence (Mr), respectively [2].

1.2.3 Magnetic Coercivity

The reverse field needed to restore magnetic induction B to zero is called coercive

field or simply coercivity (Hc). It is a good measure of the ease or difficulty of

magnetizing a material. The field needed to restore M to zero is called the intrinsic

coercivity (Hic). In soft magnetic materials, the distinction between Hc and Hic is not

important because for these materials, Hc ≪ 𝑀 so that M = 0 for essentially the same

field that gives B = 0. If the process of magnetization reversal is purely rotation, then Hc


5

is caused by the anisotropy, which resists the rotation process, and if the reversal process

is domain wall motion, then Hc is caused by the spatial variation of domain wall energy.

Often both processes determine the coercivity [3].

Definitions of these magnetic phenomena can be easy understood from the

magnetic hysteresis loops shown in the Figure 1.2 (a).

1.2.4 Magnetic Anisotropy

Magnetic properties can be anisotropic in a material. The easy axis is defined as

the easy direction of magnetization and the hard axis as the hard direction of

magnetization. Magnetic anisotropy simply means that the magnetic properties depend on

the direction in which they are measured, and this is an important property of magnetic

materials. Improvement in most magnetic devices has been made possible through the

engineering of anisotropy. Magnetic anisotropy energy is the energy that must be

supplied to work against the anisotropy force to turn the magnetization vector away from

the easy direction to the non-easy direction of magnetization. The magnetic anisotropy

energy is dependent on crystallographic directions, stress state, shape, and magnetic

history of the material. Most important anisotropies are magneto-crystalline, magneto-

static, and magneto-elastic [4].

Magnetic anisotropy energy depends upon the crystal system and direction. This

dependence is mainly due to spin-orbit coupling [5]. In Fe, which is a body centered

cubic (bcc) system, the <100> directions are easy axes of magnetization and the <111>

directions are known to be hard axes as shown in Figure 1.2 (b). The anisotropy energy,

E in a cubic single crystal system can be expressed by,

E = K0 + K1(α12α2

2 + α2

2α3

2 + α3

2α1

2) + K2(α1

2α2

2α3

2) + …….. (1.4)

where K0, K1, K2,……are anisotropy constants and are expressed in erg/cm3, and α1, α2,

and α3 are the direction cosines of the angles between the easy axes and the applied

magnetic field.

The physical shapes of the magnetic materials also contribute magnetic energy,


6

Figure 1.2 (a) A typical magnetic hysteresis loop revealing basic magnetic quantities

Figure 1.2 (b) M – H loop of bcc iron with easy and hard axis of magnetization


7

called magneto-static energy or shape anisotropy energy. The magnetization inside a

magnetized material can create magnetic free poles on surfaces of the material or at the

interface between different phases as shown in Figure 1.3. The farther these poles are

separated, the lower the magneto-static energy becomes.

The dimensions of a magnetic material change as the magnetization changes. The

resulting strain is called magnetostriction (λ), which is defined as

where l is the length measured without a magnetic field and ∆𝑙 is the length measured

with a magnetic field. The magnitude of magnetostriction becomes a useful parameter

when it is measured after materials are saturated, known as saturation magnetostriction.

The physical origin of magnetostriction comes from spin-orbit coupling, and thus the

saturation magnetostriction is an intrinsic property of the materials.

1.3 Soft Magnetic Materials

When the magnetization processes like domain wall motion and domain

magnetization rotation occur in weak fields, H < 4π Oe, which may be instantly

generated by a modest current through a few turns of wire, the material is called soft

magnet. In some very soft magnetic materials such as certain crystalline NiFe alloys

(permalloys) or amorphous metallic alloys, Hc can be as low as 12 mOe (1.0 A/m). For

comparison, the earth‟s magnetic field is about 0.4 Oe (30 A/m). Other soft materials

include pure Fe, Fe with up to 6% Si, Ni, many FeNiCo alloys, and some ferrites such as

Mn-Zn or Ni-Zn ferrites [6].

Applications of soft magnetic materials exploit the large flux changes that occur

in these materials with relatively weak changes in applied field. Soft magnetic materials

are used in transformers, inductors, motors, and generators and as field sensors in

magnetic recording or as stress/strain gauges [7].

The area inside the B-H loop, of the order of 4BrHc, is the energy per unit volume

lost per cycle in magnetizing the material. It is called the hysteresis loss. The product BH


8

Figure 1.3 Free pool distributions in magnetized bodies depending on shape anisotropy


9

has units Wb.A/m3

= J/m3. In alternating-current (AC) applications the power loss is the

frequency times the AC loop area (Jm-3

.s-1

).

1.4 Surveying High Magnetic Moment Soft Materials

1.4.1 Iron:

Among soft magnetic materials, electrical grade steel, with iron as a major

constituent, is currently being employed in the largest quantities. The annual demand of

the electronics industry amounts to several hundred of thousands of tons. The major part

of this material is used for the generation and distribution of electrical energy, of which

the application in motors takes a prominent position.

1.4.2 Fe-Si Alloys:

At the beginning of the twentieth century, it was discovered that the addition of a

few percent of Si to Fe increases the electrical resistivity and reduces the coercivity. The

latter property leads to higher permeability and lower hysteresis losses. The former

property is important because it reduces eddy-current losses. The eddy-current losses

increase with the frequency squared and can become a major problem in high-frequency

applications. This discovery led to a widespread application of Si-Fe alloys with some

compromise over saturation magnetization due to Si addition.

The random orientations of the grains in normally cast Fe-Si alloys imply that

magnetic saturation can be reached only by applying magnetic fields considerably higher

than the coercivity. This limits the useful maximum magnetic flux B to about 1.0 T. On

the other hand, hysteresis loops of single crystals are nearly rectangular so that only fields

slightly higher than the coercivity are required to drive the core to saturation. This fact

has been used in the development of grain-oriented sheets of Fe-Si with considerably

improved properties.

1.4.3 Fe-Ni Alloys:

Several magnetic alloys, for instance Ni-Fe alloys, can acquire magnetic

anisotropy when annealed below their Curie temperature. Materials having a fairly square

hysteresis loops are obtained when the annealing is performed in the presence of an

applied magnetic field. The hysteresis loop may become constricted if no field is present.


10

The anisotropy obtained in a magnetic material by annealing in a magnetic field

is called thermomagnetic anisotropy. Its occurrence is due to short range directional

ordering of atomic pairs. The magnetic-coupling energy of a pair of atoms generally

depends on the nature of atoms involved. Annealing below the Curie temperature in the

presence of an applied magnetic field tends to align the coupled pair atoms in a way that

they have their moments in the field direction, so as to minimize the free energy. Fast

cooling to a sufficiently low temperature then freezes the directional order obtained. It

leads to a uniaxial magnetic anisotropy, the easy axis of the magnetization direction lying

in the field direction. Hysteresis loops measured in the same direction are square. By

contrast, skew hysteresis loops are obtained when measuring in a direction perpendicular

to the direction of the alignment field.

An unmagnetized piece of magnetic material does not have net magnetization

because it is composed of an assembly of magnetic domains with different magnetization

directions in a way as to minimize the magneto-static energy. Non magnetized Ni-Fe

alloy bears similar type of domain pattern as well. If no field is applied during annealing

treatment, the pair moments will become aligned in the local field corresponding to the

local magnetization in each magnetic domain.

1.4.4 Fe-Al and Fe-Al-Si Alloys:

It is an important group of soft-magnetic materials that are primarily applied in

recording heads. These materials are characterized by high electrical resistivities, high

hardness, high permeability, and low magnetic losses. Optimal magnetic properties are

obtained in a fairly narrow concentration range around 9.6% Si, 5.4% Al, and 85% Fe.

This material is also known under the name Sendust.

1.4.5 Soft Ferrites:

Soft ferrites are known to have very low magnetic anisotropy. These materials can

be visualized as consisting of mixed oxides and have the general formula AB2O4 where A

and B represent various metal cations. These materials are primarily used in high

frequency applications where reduction of the various losses accompanying high

frequency magnetization is more important than the static magnetic characteristics.


11

1.4.6 Amorphous Soft Magnetic Alloys:

Some amorphous alloys have also been found to exhibit soft magnetic properties

much superior to those found in crystalline materials. For instance, the core losses

measured in amorphous alloys of the composition Fe72Co8Si5B15 have values that are

about an order of magnitude smaller than those of commercial Fe-Si alloys.

In amorphous state, the constituting atoms have a more or less random

arrangement, with grain boundaries being absent. The amorphous state is less stable than

the crystalline state and this causes amorphous alloys to spontaneously crystallize upon

heating. This amorphous-to-crystalline transformation takes place at the crystallization

temperature (Tx) which depends on the composition of the alloy. Most amorphous alloys

show a slight atomic rearrangement at temperatures somewhat below Tx, known as

structural relaxation.

1.4.7 Nanocrystalline Alloys:

These alloys have a microstructure consisting of ultrafine grains in the nanometer

range. The first step in the manufacturing of nanocrystalline alloys is the same as used for

amorphous alloys. Subsequently, these alloys are given a heat treatment above the

corresponding crystallization temperature. The composition of nanocrystalline alloys has

been slightly modified with respect to that of soft magnetic metallic glasses and contains

small additions of Cu and Nb [8].

1.5 Iron Nitride (FeN) Based High Density Head Materials

The key concerns in the research of new head materials are, how to develop high

saturation moment materials and how to decrease anisotropy energy and magneto-

striction at the same time. Amorphous materials have been studied because of the

decreased magnetocrystalline anisotropy energy that results in phase transformation.

Nanocrystalline materials and laminated systems have also shown soft magnetic

properties, with a reasonable saturation magnetic flux density. The small grains in

nanocrystalline materials and the magnetostatic coupling between the magnetic layers in

laminated systems are responsible for soft magnetism. In this section, Fe-based thin film

head materials deposited by physical vapor deposition (PVD) technique using Ar or an

Ar-N2 gas mixture are being discussed.


12

1.5.1 FeN-Based Amorphous Films:

For the amorphous films of a nitrogenated Fe system, it is found that too much

nitrogen in the amorphous Fe-N results in a coercivity of over 600 Oe, which is not

suitable for either head or medium applications [9]. Less than 10% nitrogen content in the

Fe-N amorphous film could help lower the coercivity while maintaining a reasonably

high saturation magnetization. An amorphous film deposited by dc reactive magnetron

sputtering in an Ar-N2 mixture gas has been reported, having a coercivity of ≈2 Oe, an

anisotropy field of 10 Oe, a saturation magnetostriction of 2 x 10-6

, a (saturation)

magnetic flux density of ≈18 kG, and a resistivity of ≈145 μΩ-cm [10]. Though these

amorphous films have good soft magnetic properties and a reasonable saturation

magnetic moment, these FeN-based amorphous materials have been facing a problem

regarding applications: at higher temperatures, they decompose into α-Fe and γ′-Fe4N.

1.5.2 FeN-Based Nanocrystalline Films:

Even though iron has high saturation magnetization and high permeability, it is

very difficult to induce soft magnetic properties into an Fe film because of Fe‟s large

magnetocrystalline anisotropy. It has been turned out that the incorporation of nitrogen

into an Fe matrix reduces crystalline grain size so that the average anisotropy energy of

the film becomes smaller, which leads the magnetic properties to become soft, enough for

thin film head applications. Terada, et al. [11] have demonstrated the capability of

nitrogenated Fe films fabricated by ion-beam deposition for thin film heads consisting of

α-Fe, γ′-Fe4N, and an unknown phase.

Since Fe-based nitride films possessing abnormally high saturation magnetization

were reported by Kim and Takahashi [12], close attention has been paid to Fe-based

nitride (FeXN) materials, with the aim of incorporating them into thin-film heads for high

density disk drives. The researchers attributed this giant magnetic moment to the

presence of α″-Fe16N2 phase of iron nitride. These high moment Fe-based nitride phases

will be reviewed in detail in the next chapter.

1.5.3 FeN-Based Nitride Multilayers:

Fe-based nitride multi-layers using both metallic and insulator spacer have also

been studied. For instance, γ′-Fe4N/Fe multi-layers have been deposited with improved


13

crystallinity after annealing at 200 ˚C [13]. When a 3.5 nm thick Al layer was used as a

spacer, the magnetic properties of Fe-N laminates became softer, and the saturation

magnetization increased as the thickness of the Fe-N layer decreased from 270 to 3.2 nm.

Moreover, SiO2 [14] and Si-N [15] spacer layers have also been employed to improve the

thermal stability and soft magnetic characteristics of Fe-based nitride multi-layers.

1.6 Thin Film Processing

1.6.1 Growth Process

Any thin film growth process involves three main steps:

1. Production of appropriate atomic, molecular, or ionic species.

2. Transport of these species to the substrate through appropriate medium.

3. Condensation on the substrate, either directly or via a chemical and/or

electrochemical reaction to form a solid deposit. Formation of thin films takes

place via nucleation and growth process.

1.6.2 Factors Affecting Film Properties

Deposited thin films generally have unique properties compared to the materials

in the bulk form. The specific properties of any thin films formed by any atomistic

deposition process depend on the following four factors [16].

Substrate surface condition before and after cleaning: surface morphology

(roughness, inclusions, particulate contamination), surface chemistry (surface

composition, contaminants), mechanical properties, surface flaws, out-gassing,

preferential nucleation sites, and the stability of the surface.

Details of the deposition process and system geometry: deposition process employed,

angle of incidence distribution of the depositing ad-atoms, substrate temperature,

deposition rate, gaseous contamination, concurrent energetic particle bombardment

(flux, particle mass, energy).

Details of film growth on the substrate surface: condensation and nucleation of the

arriving atoms, interface formation, interfacial flaw generation, energy input to the

growing film, surface mobility of the depositing ad-atoms, growth morphology of the


14

film, gas entrapment, reaction with deposition ambient (including reactive deposition

processes), changes in the film properties during deposition.

Post-deposition processing and reactions: chemical reaction of the film surface with

the ambient, subsequent processing, thermal or mechanical cycling, corrosion,

interfacial degradation.

Each of these factors must be reproducible in order to have reproducible film properties.

1.7 Sputtered Films and Recording Applications

Grove first observed sputtering in a dc gas discharge tube in 1852 [17]. He

discovered that when cathode surface of the discharge tube was sputtered by energetic

ions in the gas discharge, the cathode materials were deposited on the inner wall of the

discharge tube.

In a conventional sputtering system, sputtered atoms are generally composed of

neutral single atoms of the target material when the target is sputtered by bombardment

with ions having a few hundred electron volts. These sputtered atoms are partially ionized

in the discharge region of the sputtering system. This sputtering phenomenon, which was

regarded as „an undesired one‟ at the time of its discovery since it destroyed the cathode

and the grid in the discharge tube, is now widely being used for the growth of thin films

for technological aspects of research and applications.

In magnetic recording media, the films deposited by the sputtering process

produces a narrow magnetic gap for videotape recording systems and for computer disk

applications. In the production of the magnetic gap, a nonmagnetic spacer is formed from

glass material. Thin-film deposition technology enabled the production of magnetic heads

with narrow gap lengths of 0.3 μm [18]. The narrow-gap forming technology was based

on the atomic scale achievable by thin-film deposition processes. Sputtering technology,

with its precise, controlled deposition, is used to develop layered new materials including

giant magnetoresistance (GMR) magnetic materials. The spin-dependent, tunneling-

magnetoresistance (TMR) effect is expected to provide a high-density memory disk of up

to 200 Gbit/inch2 [19].


15

1.8 Brief Overview of this Dissertation

After having some essential knowledge of basic magnetic quantities and the

growth mechanism of thin films, we proceed ahead. As this dissertation focuses the

investigations on „giant‟ magnetic behavior of iron nitride, so two strong ferromagnetic

phases of this system have been selected to examine systematically.

Chapter 2 deals with the literature review of iron nitride family. The rising

interests of the researchers in γ′-phase and the exciting controversy about giant magnetic

moment of α″-phase have been discussed in separate sections. Further two sections have

been dedicated to the review of literature describing alloy addition in α″-phase and

multilayered iron nitride films.

Sample preparation methods and characterization techniques adopted during this

research are fully described in chapter 3. This chapter describes comprehensively the

magnetron sputtering system, and characterization tools such as X-ray diffraction (XRD),

scanning electron microscope (SEM), atomic force microscope (AFM), transmission

electron microscope (TEM), magnetic force microscope (MFM), magneto-optical

microscope magnetometer (MOMM), vibrating sample magnetometer (VSM), and

superconducting quantum interference device (SQUID).

The results and discussion have been divided into three chapters. The preparation

and characterization of γ′-Fe4N thin films have been described in chapter 4. The optimum

conditions to grow highly structured films have been discussed and the consequences of

this high quality crystalline growth on the saturation magnetization have been described.

The discussion about the influence of surface roughness, full width at half maximum, and

lattice mismatch on the magnetic properties of γ′-Fe4N films is also a part of this chapter.

Chapter 5 is related to the preparation and characterization of α″-Fe16N2 films.

The thermal stability is the key issue in this phase of iron nitride which is supposed to be

solved by alloying with the addition of certain metals. In this chapter, the influence of Co,

Pt, and Cr as alloy addition on the saturation magnetization of α″-phase of iron nitride is

described.


16

Chapter 6 is devoted to study the Barkhausen effect in 2D γ′-Fe4N films.

Magnetization reversal process takes place in the form of discrete and jerky jumps which

are random in terms of size, interval, and location of the jump. The answer to the

question, “Is there any law governing this seemingly random behavior?”, is also

investigated in this chapter.

The summary and conclusions from our research are presented in chapter 7.

Chapter 2 Iron Nitride and Related Compounds

17

Iron Nitride and Related Compounds

Among ferromagnetic materials, iron nitrides and related compounds have

attracted a considerable attention of researchers owing to their amazing magnetic

properties. These magnetic properties arise from the interstitial modification of 3d iron

metal by nitrogenation, leading to the expansion of the crystal lattice and changes in the

electronic structure, which in turn produces significant changes in the intrinsic magnetic

properties of these iron-based materials. Thus, they become strong ferromagnets with

increased magnetization. In this chapter, we have reviewed the attempts made by the

researchers to explore these materials since their discovery.

2.1 Iron Nitrides

Iron Nitride has been known for its variety of phases with different crystal

structures emerging from different nitrogen concentrations. The phase and the crystal

structure of iron nitride films change from α′-FexN (x ≥ 8) having tetragonal structure

[20], α″-Fe16N2 with a base centered tetragonal (bct) structure [21], γ′-Fe4N with face

centered cubic (fcc) structure [22], ε- FexN (2 < x ≤ 3) bearing hexagonal closed pack

(hcp) structure [23], to ζ-Fe2N appearing with an orthorhombic structure [22], as the

nitrogen concentration increases.

Regarding the cubic FeN phases, the ZnS-type γ″-FeN is a non-magnetic

substance in which Fe atom has no local magnetic moment, and the NaCl-type γ‴-FeN is

an antiferromagnetic that shows a large hyperfine magnetic field of up to 49 T at 4.2 K

[24] which suggest that it should be regarded as an ionic compound. FeN0.63 and FeN0.65

contain both the NaCl-type (a = 0.450 nm) and ZnS-type (a = 0.433) nitrides, while

FeN0.91 is nearly single phase of ZnS-type nitride. These two nitrogen-rich phases are

relatively new and less well known.


18

As the concentration of N is lowered, the iron sub-lattice transforms into a

hexagonal phase called ε-phase, that can accept nitrogen in the concentration range from

25 at.% up to 33 at.%. The atoms are randomly distributed in the octahedra formed by the

Fe sub-lattice. At the Fe2N stoichiometry, the ε phase can be transformed into ζ-Fe2N,

which has an orthorhombic structure, by ordering the N-atoms over the octahedral sites.

The magnetic properties also change due to the change in nitrogen concentration. The

saturation magnetization of ε- FexN (2 < x ≤ 3) phase decreases as „x‟ decreases, and it

disappears at room temperature below an „x‟ of 2.3 [25]. The ζ-Fe2N phase does not show

ferromagnetism at room temperature and the Curie temperature fluctuates around 7 K

[26].

A non-crystalline FeN phase has been produced by ball milling γ′-Fe4N with α-Fe

[27]. The phase exhibited ferromagnetic behavior with a reduced magnetization and it

was suggested that it was formed in the grain boundaries of nano-crystalline α-FeN.

Prolonged aging may lead the nano-composite to the equilibrium product α-Fe plus γ′-

Fe4N.

All phases of iron nitride are meta-stable with respect to decomposition into N2

gas and α-Fe, although the kinetics of this process is quite slow at temperatures below

670 K. However, the actual decomposition temperature depends on the phase.

In general, when x > 2, the FexN compounds exhibit ferromagnetic properties at

room temperature. As the present dissertation focuses on ferromagnetic behavior of iron

nitride, so the ferromagnetic iron nitride phases will be reviewed in next sections.

2.2 γ′-Fe4N

In magnetism, maximum interest has been focused on iron and iron-based alloys

because these materials are of greatest practical use. In addition, iron is by far the

cheapest magnetic element. Interstitial nitrogen has the potential to alter the properties of

iron-based alloys in a useful sense [28]. Due to this reason, iron nitride alloys have been

attracting researchers from worldwide since several decades.

The simplest and the most ordered phase of Fe-N system is cubic γ′-Fe4N, as

shown in Figure 2.1. The fcc lattice is expanded by 33 % by interstitial nitrogen in the 1b

body-centre position, giving a lattice parameter of 0.3795 nm, compared with 0.3450 nm


19

for fcc γ-Fe. The compound has two nonequivalent iron sites: the body-corner 1a site (Fe

I) with cubic symmetry having 12 iron neighbors at 0.268 nm, and the 3c face-centre site

(Fe II) with two nitrogen neighbors at 0.190 nm and 12 iron second neighbors at

0.268 nm [28], as listed in table 2.1.

The nitrogen induced expansion stabilizes a ferromagnetic state for γ′-Fe4N with a

Curie temperature of 767 K, and a high value of room temperature polarization.

In order to explore the intrinsic magnetic moment for γ′-Fe4N films, numerous

electronic structure calculations have been carried out. For instance, Sakuma [29,30] has

performed the self-consistent spin-polarized band calculations for γ′-Fe4N and studied the

magnetic structure combined with the electronic features. The calculations were carried

out using the linearized muffin-tin orbital (LMTO) method [31], within the frame of local

spin density functional formalism. The obtained magnetic moment of each Fe atom site

show fair agreement with the experimental measurement.

The interstitial insertion of nitrogen into iron results in (1) the change in 3d

density of states, and reduction in difference in occupancy of the 3d↑ and ↓ states due to

hybridization with the sp orbitals of nitrogen, (2) the expansion of iron lattice, reducing

the 3d-3d overlap, and hence the bandwidth, which tends to make iron a strong

ferromagnet, and (3) the change in the symmetry of iron sites. The reason why the

moments on iron sites which have nitrogen nearest-neighbors are reduced, and the

second-neighbor iron moments are enhanced is explained in Figure 2.2. The

hybridization of 3d states of the iron that is a nearest neighbor with the sp states nitrogen

reduces the spin splitting, and especially lowers the potential for 3d↓ electrons. This

results in the charge transfer from the more distant iron, predominantly from 3d↓ to 3d↓

states. This depletes the 3d↓ band of the strongly ferromagnetic distant neighbors and

hence increases their magnetic moment [28].

The value of average spin magnetic moment of γ′-Fe4N has also been observed in

close agreement as determined from full-potential band structure calculations performed

by Coehoorn et al. [32].


20

Figure 2.1 Crystal Structure of γ′-Fe4N

Table 2.1 Iron sites in γ′-Fe4N and their nearest neighbors with distances

Iron Site No. of Fe

Neighbors

Distance from

Fe Neighbors

No of Nitrogen

Neighbors

Distance from Nitrogen

Neighbors

Fe I

12 0.268 nm - -

Fe II

12 0.268 nm 2 0.190 nm


21

Figure 2.2 Schematic changes in the electronic structure at first and second neighbor

iron site, induced by hybridization with nitrogen


22

Yamaguchi et al. [33] have synthesized γ′-Fe4N both in particle and thin film

forms by nitriding iron in ammonia-hydrogen gas at 500 ˚C. Single phase γ′-Fe4N was

obtained at ammonia flow rates between 40% to 60% and exhibited saturation

magnetization of 188±2 emu/g. The films prepared at different ammonia flow rates

exhibited different crystalline orientations. In 1997, structural, magnetic, and

superconducting properties of ferromagnetic/superconducting multi-layers of γ′-

Fe4N/NbN were explored by Mattson et al. [34]. The superconducting properties showed

that 11 Å ferromagnetic layers were sufficient to decouple the superconducting layers and

to yield anisotropic behavior. On the other hand, Borsa et al. [35] studied Cu3N/ γ′-Fe4N

bi-layers and explored that a quite favorable match between Cu3N, an insulator and γ′-

Fe4N, a ferromagnetic conductor, might open the prospects of developing an all-nitride

all-epitaxial magnetic tunnel junction [MTJs]. Prior to this, epitaxial thin films of γ′-Fe4N

were grown on MgO(001) substrates [36] revealing layer by layer growth mode. The

magnetic domains showed a complete in-plane magnetization with easy axis along the

[001] crystallographic axis. The channeling yield in γ′-Fe4N films grown on MgO(100)

substrates amounted to 11%, which was close to the calculated value of 8% [37]. These

properties made γ′-Fe4N films an interesting candidate for device applications. Here, note

that the word „epitaxy‟ refers to a growth mode of thin films in which single crystal films

of a material are deposited on a single crystal substrate such that both the deposit and the

substrate have the same structural orientation.

Epitaxial γ′-Fe4N films were also grown on MgO(100) substrate [38,39] using a

halide source and the investigations implied that the epitaxial layer of γ′-Fe4N showed an

anomalous light reflectivity modulated by an external magnetic field. The phenomenon of

light reflectivity was further investigated by exploring the temperature dependent

behavior of magnetic thin films [40]. This research group also worked on the preparation

of iron nitride substrates and succeeded to fabricate a free standing γ′-Fe4N substrate by

atmospheric pressure halide vapor phase epitaxy (AP-HVPE) using FeCl3 and NH3 [41].

The thickness of the iron nitride deposited as substrate was in the range of 50-100 μm.

Loloee et al. [42] deposited NbN and γ′-Fe4N films and NbN/γ′-Fe4N bi-layers

using a dc triode magnetron sputtering in N2 reactive gas. All the films showed epitaxial

behavior and the magnetic characterization revealed that there is a strain-induced in-plane


23

uniaxial anisotropy in γ′-Fe4N films with preferred easy axis direction along <010> and

<001>, regardless of their being grown on a NbN film or a MgO substrate.

A less mismatch between substrate and the deposit might provide a better

template for epitaxial growth, so Nikolaev et al. [43] deposited γ′- iron nitride films on

SrTiO3(001). The epitaxial quality and smoothness of the films implied that the material

was well suited for spin-dependent investigations and the devices based on the

hetrostructures.

Costa-Kramer et al. [44] have grown γ′-Fe4N thin films by Molecular Beam

Epitaxy (MBE) of iron in the presence of nitrogen obtained from a radio-frequency (rf)

atomic source and studied in detail, the magnetization reversal process of pure epitaxial

γ′-Fe4N films by Kerr measurements. In order to investigate the exact growth mechanism,

involving the reaction of a gas and a metal, Gallego et al. [45] have deposited iron nitride

films on Cu(100) substrate. Single phase epitaxial films were magnetic, grown layer by

layer at room temperature, with easy axis in the plane of the film and parallel to the

<100> direction.

Substrate Materials always play an important role affecting properties of thin

films. Wang et al. [46] have reported the influence of substrate material on the structure

and magnetic properties of iron nitride thin films. Single phase γ′-Fe4N films were grown

on three substrates: Si(100), NaCl(100), and glass. The films deposited on glass

substrates had small grain size as compared to other two substrate materials and

consequently less coercivity. The influence of substrate temperature on the magnetic

properties of γ′-Fe4N films deposited on NaCl(100) have also been investigated [47].

Ecija et al. [48] have reported that the magnetization reversal behavior of γ′-Fe4N

films deposited on Cu(100) substrates depended on the angle between the field and

anisotropy axes. The polar plot of the remanence displayed a „butterfly‟ behavior, usually

not expected from simple cubic structures. Uniaxial anisotropy might be an important

factor in order to explain this phenomenon.

Turning to the application of iron nitride thin films in device fabrication, Sunaga

et al. [49] have employed γ′-Fe4N thin films in MgO-based MTJs with Fe-N electrode. A

relatively large inverse tunnel magneto-resistance (TMR) effect and a strong asymmetric


24

VB dependence were observed in TMR ratio and tunnel conductance. This is believed to

be an important progress in the technological aspects of MTJs.

However, to our knowledge, no previous attempt is reported, exploring the

annealing time and lattice mismatch dependent magnetic behavior of γ΄-Fe4N thin films.

Therefore, we have investigated the epitaxy-dependent saturation magnetization achieved

by optimizing annealing time in γ΄-Fe4N thin films and the subsequent effect of lattice

mismatch on saturation magnetization using three different substrates MgO(100),

SrTiO3(100), and LaAlO3(100), with lattice mismatch of 11%, 3%, and 0%, respectively

[50]. The static and dynamic magnetic domain patterns of thin films deposited at 200-

500 ˚C substrate temperature have also been discussed [51].The details of these

investigations will be discussed later.

2.3 α″-Fe16N2

α″-Fe16N2 has remained one of the most fantastic and mysterious materials in the

field of magnetism since the last 35 years due to its controversial magnetic moment.

Prolonged annealing of the α′-phase at about 370-420 K leads to an ordering of nitrogen

to produce the famous ordered martensite α″-phase, first described by Jack in the early

50‟s [21]. The unit cell consists of base-centre tetragonal (bct) structure with a = 0.572

nm and c = 0.629 nm, as illustrated in Figure 3.

There are three non-equivalent sites: 4d with no nitrogen neighbors, 4e with one

nitrogen neighbor at 0.193 nm, and 8h with two nitrogen neighbors at 0.179 nm.

Prolonged annealing or higher temperature treatment leads to decomposition of α″-phase

into α-Fe and γ′-Fe4N [28]. The limited solubility of nitrogen in γ-Fe, at most 10.5 %,

means that it is impossible to obtain a pure 16 : 2 phase in bulk form by the „quench and

anneal‟ route. However, in thin film form, much larger amounts of α″-phase (≥ 80%) can

be stabilized [52].

Interest in magnetic properties of α″-phase dates from the work by Kim and

Takahashi [12] on iron films evaporated in nitrogen. Their films were a mixture of α″-

phase and α-Fe, and the polarization of α″-Fe16N2 films was inferred to be as high as 2.8

T, corresponding to an average iron moment of 3.0 μB. This report made a strong impact


25

Fig. 2.3 Crystal Structure of α″-Fe16N2


26

on the research fields of magnetic materials and since then, many theoretical and

experimental attempts have been made to investigate this effect literally called „giant‟

magnetic moment.

The research on α″-Fe16N2 thin films was stimulated by the reports of very large

magnetic moments by Sugita et al. [53,54] for single crystal thin films grown epitaxially

on GaAs or (Ga0.8In0.2)As by MBE. The Bs of those films were confirmed to be 2.8-3.0 T

at room temperature, equivalent to the average magnetic moment per Fe atom of 3.1 - 3.3

μB. Beyond 400 ˚C, α″-Fe16N2 films were found to dissociate into α-Fe and γ′-Fe4N. Later

on, the magnetic moments were found to be as high as 3.5 μB/Fe atom, with a

corresponding polarization of 3.2 T [55], while the film thickness was about 50 nm,

having nitrogen concentration nearly 11%. The temperature dependence of the saturation

magnetization was found to obey T3/2

law at low temperature. The electrical resistivity

was around 30 μΩ cm and decreased linearly with decreasing the temperature showing

metallic behavior. The saturation anomalous Hall resistivity was found to be 4 x 10-7

Vcm/A, which was much higher than the values for Fe-N martensite (1.9 x 10-7

Vcm/A)

and than that of Fe (1.5 x 10-7

Vcm/A). These results were consistent with Sakuma‟s

model [56]. In a separate study of this group, the value of g factor for α″-Fe16N2 films

was reported to be about 2.0, meaning that magnetic moment mainly originates from

spin. Torque magnetometery showed a large perpendicular anisotropy of 7.8 x 106 erg/cc,

which can originate from its bct structure [57]. The value of 4πMs has been found to

increase with increase in N atom ordering by annealing and finally reaching 29 kG for all

the α″-Fe16N2 films prepared in the thickness range of 200-900 Å [58].

The controversy about giant magnetic moment in α″-phase started when some

groups could not reproduce the results presented by Kim and Takahashi [12], and

confirmed by Sugita [53-58]. For instance, Takahashi et al. [59,60] adopted two different

routes to investigate α″-Fe16N2 films systematically. One was the conventional sputtering

and the other was plasma evaporation. The value of saturation magnetization for the films

fabricated by plasma evaporation was measured to be 235 emu/g at 300 K (≈ 2.4 μB/Fe

atom on average) and for the films made by sputtering, a value of 218 emu/g was noted.

A low value of saturation magnetization was also observed in the single phase α″-

Fe16N2 films grown epitaxially on Al2O3 substrates using the thermalized plasma dc


27

sputtering process. The thermalized sputtered pressure was favorable for the growth of

α″-phase [61].

Min [62] has carried out the total energy electronic structure calculations in order

to understand the microscopic origin of the enhancement of the Fe magnetic moments in

α″-Fe16N2 films. The results revealed that spin magnetic moments for three types of Fe:

Fe-I, Fe-II, and Fe-III were 2.13, 2.50, and 2.85 μB, respectively, and the orbital

contribution was minor. Thus, large enhancement was found in Fe-II and Fe-III, which

were located far from the N atom. The observations imply that local environment plays

an important role in determining the magnetic moment of Fe atoms in iron nitride

compounds. On the other hand, Full-potential band structure calculations revealed that Fe

magnetic moment could not be more that 2.37 μB/Fe atom in α″-phase of iron nitride [32].

However, a significantly enhanced moment was observed by Huang et al. [63] in

iron nitride prepared by treating Fe powder with NH3/H2 gas mixtures at a temperature of

about 665 ˚C, followed by quenching. The α″-phase was achieved by treating α′-phase for

about 1 to 2 hours at 120-150 ˚C. Almost the same results were achieved by Bao et al.

[64] by following the same preparation route, except that the starting material was γ-

Fe2O3, which was first reduced to α-Fe under H2. Higher value of saturation

magnetization as compared to pure iron was also achieved by depositing iron nitride

films on Ge(100) wafers by a reactive ion beam sputtering method in an ammonia

atmosphere [65].

Brewer et al. [66] deposited an Ag under-layer on Si(100) substrates prior to

deposition by reactive sputtering in nitrogen. The films were a mixture of α′-phase (54 %)

and α″-phase (46 %), and showed an average moment of 1780 emu/cc, considerably

larger than that of pure α-Fe. Okamoto et al. [67] have studied the effect of lattice

distortion on magnetization and found that the average magnetic moment gradually

increased with the unit cell volume and reached up to 2.8 μB/Fe when the unit cell volume

was around 205.8 Å3 reported by jack [21].

Although, the investigations by Coey et al. [68,69] could only explore the

magnetic moment in α″-phase in the range of 2.3-2.6 μB/Fe, nevertheless, they have

encouraged further investigations, particularly involving alloy additions with a view to


28

optimize the properties for applications which call for materials with greatest possible

magnetic polarization.

We have achieved the optimum conditions, depending upon nitrogen partial

pressure, substrate temperature and annealing time, for the single phase epitaxial growth

of α″-phase of iron nitride deposited on single crystal Si(100) substrates [70]. The details

will be discussed in chapter 5 of this dissertation.

2.4 α″-(Fe,X)16N2

Proper addition of substitutional or interstitial alloying elements might favor the

stabilization of 16:2 nitride phase and enhance the value of magnetic moment [69,71].

Inoue et al. [72] have investigated the effect of Co addition in α″-phase. The reasons for

Co selection are twofold: the first point is that Fe maintains a bcc structure even after

alloying with the third element; the second point is that the third element does not

undergo preferential nitridation. Co meets the above two requirements. The results

revealed that at a formation temperature of 500 ˚C, which was 50 ˚C higher than pure

iron, the amount of 16:2 nitride phase was 3 times higher than that for pure iron,

suggesting that Co stabilizes the 16:2 nitride structure.

Takahashi et al. [73] have expanded the Fe-N system to Fe-H-N, in addition to

Fe-Co-N. Fe-H martensite was synthesized by using sputtering under Ar+H2 atmosphere.

The value of saturation magnetization of the α″-(Fe100-x,Cox)16N2 (X = 10-30) could not

exceed more than 240 emu/g (≈ 2.4 μB/Fe atom) at 300 K. On the other hand, an

enhanced value of magnetic moment has been reported by Jiang et al. [74] for Co

addition in α″-phase. Their results showed that the films with 5-25 at.% Co contents

deposited on Si(100) and NaCl(100) substrates at 100 ˚C exhibited a saturation

magnetization up to the range of 2.5-2.7 T, which was close to the value expected from

α″-alloy phases.

The C atom also enjoys the reputation of making interstitial compounds with bcc

Fe. Therefore, Takahashi et al. [75] have studied the α″-Fe16(N,C)2 system to explore the

possibility of giant magnetic moment. It was observed that the formation of α″-alloy


29

phase became more difficult with increasing C contents and the saturation magnetization

decreased as well.

Much improved thermal stability up to 500 ˚C was observed in Fe-Ti-N system by

Wang et al. [76-78]. In addition, a high value of saturation magnetization (≈24 kG), low

coercivity (≈1.5 Oe), high relative permeability (μ≈3200 at 1 MHz), and high electrical

resistivity (≈ 100 μΩcm ) were measured. All these properties herald promise for

applications in high density recording. Byeon, et al. [79] have also confirmed the

increased stability of α″-Fe16N2 system by Ti addition but could not achieve a high value

of saturation magnetization as reported ealier [76-78].

We have investigated the effect of Co, Pt, and Cr additions on the saturation

magnetization of α″-phase of iron nitride [70]. The details will be discussed in chapter 5.

2.5 Polycrystalline/Multiphase Iron Nitrides

As mentioned earlier, research on iron nitride has a history spreading over quite

some decades. During the period, many research groups could not prepare any phase of

iron nitride thin films in pure form. The films mostly consisted of multi-phase iron

nitrides or a slight appearance of single phase which would decompose into other

phase/phases by a quite small variation in deposition parameters. For example, the films

prepared by Xiao et al. [80] using a mixture of NH3 and Ar gases by reactive sputtering,

comprised ζ-Fe2N, ε- Fe2-3N, ε- Fe3N, and γ′-Fe4N phases.

The films prepared by Guibin et al. [81] using ion beam enhanced deposition

(IBED) method, consisted of α-Fe, α″-Fe16N2, γ′-Fe4N, and ε- Fe2-3N phases or a mixture

of these phases.

Multi-phase iron nitride films with a relative high magnetic moment than iron

have also been reported by a reactive ion beam sputtering method on Ge(100) wafers

[82], and by radio-frequency magnetron sputter deposition on Si (100) substrates [83,84].

Effect of nitrogen concentration on the Fe-N films has been investigated [85,86],

and a high value of saturation magnetization (≈2223 emu/cc) and low coercivity (≈4 Oe)

has been found at rf power of 200 W and nitrogen partial pressure of 2 x 10-4

Torr [86].


30

Iron nitride coated iron micro-particles were prepared by Luo et al. [87] by in-situ

surface nitridation of iron particles with NH3 gas at a temperature of 510 ˚C. The particles

exhibited greatly improved chemical stability, while maintaining a large magnetization,

which makes them interesting for applications in water-based magneto-rheological fluids

and polishing fluids.

Chapter 3 Experimental Techniques

31

Experimental Techniques

Magnetic thin films deposited by magnetron sputtering enjoy many advantageous

features, so this technique has been employed in the present research work. The films

were characterized by X-ray diffraction, atomic force microscopy, scanning electron

microscopy and transmission electron microscopy for its structural properties. The

magnetic properties have been explored by using vibrating sample magnetometery,

magnetic force microscopy and magneto-optic microscope magnetometer. A

superconducting quantum interference device was employed to investigate the magnetic

and electrical behavior of the films at low temperatures. The compositional analysis

performed using Auger electron spectroscopy has also been described in this chapter.

3.1 Magnetron Sputtering

3.1.1 The Sputtering Principle

In a sputtering process, gas ions from a plasma are accelerated towards a target

consisting of the material to be deposited. The material is detached or „sputtered‟ from

the target and subsequently deposited on a substrate placed in the vicinity. The process is

accomplished in a closed recipient, which is pumped down to a vacuum base pressure

before the start of deposition.


32

3.1.2 The Sputter Parameters

The sputtered film properties are controlled by adjusting the following sputter

parameters.

The sputter current Isp determines mainly the deposition rate and hence the time which

the particles take during the growth process for either, surface diffusion and

agglomeration on existing growth centers or for the nucleation with other adatoms.

The applied voltage determines the sputter yield and the maximum energy, with which

the sputtered particles can escape from the target. The sputtered particles show a broad

distribution of energies.

The pressure P in the sputter chamber determines the mean free path λ for the sputtered

material, which is proportional to the 1/P.

The Target to substrate (TS) distance together with the pressure controls how many

collisions of the particles occur on their way from the target to the substrate. This can

affect the porosity, crystallinity, and the texture of the film [88].

The substrate temperature can have a strong effect on the growth behavior with respect to

the crystallinity and/or density of the samples. It can be adjusted from the room

temperature up to 500 ˚C. However, during sputtering even without external heating, the

substrate temperature may rise considerably for long sputtering times required for the

deposition of thick films [89].

A bias voltage can be optionally applied to the substrate up to ±100 V, which has the

effect of accelerating electrons or ions towards the substrate or keeping them away. Both

may have an influence on the layer growth [90,91].

3.1.3 The Sputter System

The sputter deposition of magnetic thin films is quite common as it offers the

advantages of high deposition rate, large sample throughput, low cost, and stoichiometric

equality of sputter target and the grown layer as well. A conventional dc magnetron

sputtering system is shown in Figure 3.1. A heavy inert gas, most frequently Ar, is bled


33

Figure 3.1 Schematic diagram of a dc magnetron sputtering system


34

into the chamber at a pressure in the range 10-3

to 10-1

Torr during growth. The base

pressure in the growth chamber prior to deposition is commonly more or less 10-8

Torr.

The films are then sputtered from bulk targets placed at some distance from the

substrates. During the sputter process, the electrons that are ejected from the cathode, are

accelerated away from the cathode and are not efficiently used for sustaining the

discharge. By the suitable application of a magnetic field, the electrons can be deflected

to stay near the target surface and by an appropriate arrangement of the magnets, the

electrons can be made to circulate on a closed path on the target surface. This high flux of

electrons creates high density plasma from which ions can be extracted to sputter the

target material producing a magnetron sputtering configuration [92].

Two types of sputter depositions are commonly employed depending upon the

nature of power source [93]. DC sputtering is usually employed for metal deposition, as

in our case, and happens when a DC voltage is applied to the target relative to the

substrate. This voltage excites plasma between target and substrate. Positively charged Ar

ions knock off atoms from the negatively charged target through bombardment, and some

atoms from the target reach the substrate forming the thin film deposit. RF sputtering can

be used for all types of materials including conductive and non-conductive materials and

uses an AC power source with RF frequency. Despite the fact that most sputter

depositions use a single gas of Ar, some times a second gas species, such as N2, is

intentionally mixed into the sputtering gas. The second gas is usually more reactive than

Ar, and as a result the element of the second gas appears in the film.

This technique is called „reactive sputtering‟, and was used in our work to deposit

iron nitride films by introducing N2 gas in addition to Ar during deposition.

3.2 Sample Preparation

3.2.1 Cleaning of Substrates

The procedure of substrate cleaning is quite essential and demands special care in

order to fabricate samples with desired composition and crystal quality. The key purpose

of cleaning the substrates is to remove the contaminations from the surface and to prepare

a defect free atomically flat surface of the substrates. Depending upon the physical and

chemical nature of the substrate material, different cleaning processes are employed. For


35

growth of γ′-Fe4N films, we have used MgO(100), SrTiO3(100), and LaAlO3(100)

substrates. Each of these substrates requires a unique cleaning recipe.

Having a rock salt cubic structure, crystals of MgO(100) are widely used as

substrates due to their chemical inertness and relatively simple cleaning process. The

organic contaminants were removed by rinsing the substrates in organic solvents like

acetone and ethanol in ultrasonic bath at 60 ˚C. Then the substrates were set on the Mo

holder and inserted into the growth chamber for annealing at 800 ˚C in O2 atmosphere

for 1 hr [94]. By employing this treatment, carbon and oxygen containing impurities were

oxidized and removed.

SrTiO3(100) is an insulator with a perovskite crystal structure where planes of Sr

mono-oxide and Ti dioxide are placed alternatively parallel to the <100> plane of the

crystal. The initial cleaning was performed in the same way as for MgO: ultrasonic

rinsing using acetone followed by ethanol to remove organic contaminants. Finally the

substrates were annealed in O2 atmosphere at 600 ˚C for one hour to remove probable C

and/or O-containing contaminants [95].

LaAlO3(100) has a rhombohedral structure at room temperature, but can be

regarded as a pseudocubic with lattice constant of 3.79 Å, perfectly as that of our desired

γ′-Fe4N films. Though relatively costly, it is emerging as a widely used substrate for thin

film growth. The initial cleaning of this substrate was carried out in boiling acetone and

isopropyl alcohol [96]. The process was followed by heating in vacuum at 450 ˚C for

30 minutes.

The Si(100) substrate surface characteristics including contaminations on the

surface affect the uniformity of the films. Exposure of Si wafers to air introduces

agglomeration of the dust particles on the surface even when the processing is done in a

clean room, so this effect should be minimized by adopting appropriate cleaning method.

The standard organic cleaning solvents like methanol, ethanol, isopropanol, and acetone

can introduce damage to the Si surface and therefore, should be avoided [97]

RCA 2 [6 parts deionized (DI) water, 1 part 27% HCl, 1 part 30% H2O2], which is

one of the most reliable methods for Si substrate cleaning, was employed. RCA 2

solution is prepared by mixing DI water and HCl, followed by heating the solution to


36

70 ˚C on a hot plate. After removing from the heat, H2O2 is added in small increments

[94]. Initially, Si substrates were cleaned using DI water and then immersed in RCA 2

solution for 10 minutes. Finally, the substrates were again rinsed with DI water. Note that

here RCA stands for „Radio Corporation of America‟ which introduced this method for

cleaning Si wafers for its semiconductor industry.

3.2.2 Preparation of γ′-Fe4N Thin Films

For the growth of single phase epitaxial films, proper selection of substrates is the

most important parameter as it is expected that the minimum lattice mismatch will

guaranty maximum epitaxy. Lattice mismatch is defined as the percentage difference of

the lattice constants between the deposited film and the substrate. Although MgO(100)

has 11% lattice mismatch with γ′-Fe4N and can be successfully employed for epitaxial

growth, as in our case too, but such a large mismatch could affect the growth normals and

the in-plane orientation relationship between the substrate and the deposit [42].

Therefore, in this research, we have employed three different single crystal substrates,

MgO(100), SrTiO3(100), and LaAlO3(100) having 11%, 3%, and 0% lattice mismatches,

respectively with γ′-Fe4N films. The main purpose of this choice was to investigate the

effect of lattice mismatch on the epitaxy of the γ′-Fe4N films and its subsequent effect on

the saturation magnetization of the films. Initially, several series of iron nitride thin films

were deposited on MgO(100) to obtain the optimized conditions for epitaxial growth as

the other two substrates were relatively costly.

Substrates were cleaned ultrasonically using acetone and ethyl alcohol, and pre-

heated for 30 minutes in vacuum at the deposition temperature prior to the deposition.

Highly pure (99.95%) target of α-iron (diameter = 50 mm) was placed at a distance of 10

cm from the substrate holder. The base pressure of the chamber was better than 2 × 10-6

Torr. Optimized partial pressures of 5 and 0.5 mTorr of analytically pure Ar (99.95%

purity) and N2 (99.95% purity), respectively, with a total working pressure of 5.5 mTorr

were injected into the chamber utilizing dc sputtering power of 30 W. A series of films

having 550 Å thicknesses was deposited at a deposition rate of 55 Å/sec at the substrate

temperature of 450 ºC and in-situ annealed for 10 to 40 minutes. Another series with

500 Å thicknesses for each film was deposited at the substrate temperatures ranging from

200 to 500 ºC and all the samples were in-situ annealed for 30 minutes.


37

3.2.3 Preparation of α″-Fe16N2 Thin Films

Single phase epitaxial growth of α″-Fe16N2 thin films was achieved by varying the

nitrogen partial pressure using single crystal Si(100) substrates. The substrates were

cleaned ultrasonically and pre-heated in vacuum at deposition temperature for

30 minutes. Analytically pure target of α-iron was placed at a distance of 10 cm from the

substrate holder. The base pressure of the chamber was better than 2 × 10-6

Torr. Ar gas

(99.95% pure) was used as a working gas and injected into the chamber at a working

pressure of 5 mTorr using dc sputtering power of 60 W. Partial pressure of nitrogen

(99.95% pure) was varied from 0.2 to 1.0 mTorr. 400 Å thick films were deposited with a

deposition rate of 0.88 Å/sec at the substrate temperature of 200 ºC and all the films were

in-situ annealed for 1 hour.

3.3 Crystal Phase Determination

3.3.1 X-ray Diffraction

The investigation of crystal structure, pole figure and phi scan were carried out

using Rigaku D/MAX-RC MPA x-ray diffractmeter (XRD) utilizing Cu Kα radiation

generated at 60 kV and 100 mA. The measurements were recorded within the 2θ range of

20˚ to 70º at a scanning rate of 2º/min. The obtained data was analyzed using MDI/

JADE5 software. XRD is an important non-destructive technique commonly utilized for

crystal structure determination. The pattern produced by the diffracted beam of x-rays

through the closely spaced lattices in a crystal is obtained and analyzed. This generally

leads to the understanding of crystal structure [98].

A crystal structure is a regular three dimensional distribution of atoms in space.

These are arranged to form a series of parallel planes separated from one another by a

distance d, which varies according to the nature of the material. When a monochromatic

x-ray beam with a specific wavelength (λ) is projected on a crystallographic material at

an angle (θ), diffraction occurs only when the distance travelled by the reflected x-rays

from the successive planes differs by a complete number of wavelengths. This geometry

in the crystal lattice, as shown in Figure 3.2, is explained using Bragg‟s law:


38

Figure 3.2 Basic features of a typical x-ray diffraction experiment


39

nλ = 2d sin θ … … … … … (3.1)

By varying the angle, the conditions of Bragg‟s law are satisfied by different d

spacings in the crystalline materials. Plotting the angular positions and intensities of the

resultant diffracted beams of radiation produces a pattern, which reveals the unique

characteristics of the crystal structure of the sample.

3.3.2 Rocking Curve

An X-ray rocking curve (RC) measurement using x-ray diffraction is widely used

for analyzing the preferred orientation of highly oriented epitaxial thin films. The rocking

curve is defined as a peak profile of diffraction which is illustrated in a graph with the

angle of incidence of X-rays on the abscissa and an intensity of the diffracted X-rays on

the ordinate. RC analysis is also known as ω-scan. Such a scan is obtained by „rocking‟

the film surface about an axis perpendicular to the plane containing the incident and

diffracted beams, while the detector position remains constant. As illustrated in Figure

3.3, ω is angle between sample surface and the incident ray, while 2θ is the angle

between incident ray and the detector. The measurement is made by doing a ω-scan at a

fixed 2θ angle. The value of full-width at half maximum (FWHM) of the RC is used as a

measure of degree of the preferred orientations [99].

3.3.3 Phi Scan and Pole Figure

The phi-scan is a very useful technique which is employed to investigate the in-

plane crystallographic relationship among crystal planes of the film deposited and

between the film and its substrate [100]. Therefore in order to examine, whether epitaxial

growth has occurred, and to determine the in-plane orientation of γ′-Fe4N films with

respect to the major axis of the MgO substrate, XRD phi-scan was carried out. For the

scan of (200) reflection of γ′-Fe4N, the 2θ was fixed at 47.91º and ω at 23.96º [101]

(where ω is the half of corresponding 2θ value).

A pole figure is basically a graphical representation of the orientation of objects in

space. In crystallography, the pole figures are used for texture analysis in the form of

stereographic projections to represent the orientation distribution of crystallographic

lattice planes in thin films [99]. The pole figures are obtained by rotating the sample


40

Figure 3.3 Definitions of angles in a rocking curve measurement set up


41

around a φ axis from 0˚ to 360˚ in steps of 1˚, again at fixed values of 2θ and ω. Here, φ

is an axis of in-plane sample rotation. In our study, the in-plane orientation of the γ′-Fe4N

film on the MgO substrate was also investigated by pole figure measurements for the

(200) reflection.

3.4 Atomic Force Microscopy

Film thickness calibrations and characterization of surface morphology were

performed using a PSIA, XE-100 atomic force microscope (AFM). The working of the

AFM is in Figure 3.4.

The AFM was conceived as a response to the question: if surfaces could be

imaged by a current, why not by a force? A major advantage of detecting forces rather

than current is that all kinds of material surfaces including metals, semiconductors, and

insulators can be imaged.

In order to mechanically sense the atomic scale surface topography using AFM, a sharp

tip is mounted at the end of the soft cantilever spring having a spring constant smaller

than the spring constant that effectively exists between atoms. Furthermore, the applied

force should not be large enough to displace the surface atoms. Normally, the AFM

cantilevers are micro-machined from silicon, silicon oxide, or silicon nitride, and have a

spring constant ranging from 0.1 to ≈ 50 N/m.

AFM has two basic modes of operation. When the tip and the specimen are

widely separated, Van der Waals forces cause them to attract weakly. But when they are

drawn too closely together, their electron clouds overlap and electrostatic repulsive forces

physically push them apart. In former non-contact AFM (NC-AFM) mode, the cantilever

is located tens to hundreds of angstroms from the specimen surface [103]. To prevent

surface contact, a soft cantilever is used resulting in low tip-specimen forces of ≈ 10-12

N/m. To detect small forces, the cantilever vibrates and changes in vibrating amplitude

(due to topography) are detected by sensitive AC methods, converted to tip-sample

spacing, and ultimately recorded as surface images [104].


42

Figure 3.4 Schematic drawing of an atomic force microscope


43

In the latter mode, called the contact mode of AFM, the tip actually makes

physical contact with the surface and forces in the range of 10-6

to 10-8

N are typically

generated.

3.5 Scanning Electron Microscopy

A Hitachi S-4800 scanning electron microscope (SEM) was also employed to

study the surface morphology and perform thickness calibration through cross-sectional

SEM-images of iron nitride thin films. As „seeing is believing‟ and understanding, the

SEM is perhaps the most widely used thin film characterization instrument. In the SEM,

however, only a small portion of the total image builds up serially by scanning the probe.

A schematic of the typical SEM is shown in the Figure 3.5. Electrons

thermionically emitted from a tungsten or LaB6-cathode filament are drawn to an anode

and focused by two successive condenser lenses into a beam with a very fine spot size

that is typically 10 Å in diameter. Pairs of scanning coils located at the objective lens

deflect the beam either linearly or in raster fashion over a rectangular area of the

specimen surface. Electron energies having ranges from a few keV to 50 keV, are utilized

[105]. Upon impinging on the specimen surface, the primary electrons decelerate and in

the process the lost energy is transferred inelastically to other atomic electrons and to the

lattice. Through continuous random scattering events, the primary beam effectively

spreads and fills a teardrop-shaped interaction volume with a multitude of electronic

excitations. The result is a distribution of electrons which manage to leave the specimen

with an energy spectrum. In addition, target x-rays are emitted and other signals such as

light, heat, and specimen current are produced; the sources of their origin can be imaged

with appropriate detectors [106].

3.6 Magnetic Force Microscopy

The static magnetic domain structure was investigated by magnetic force microscope

(MFM) which was a non-contact force microscope (PSIA, XE-100) equipped with a

magnetic tip (Nanosensors). MFM is a very powerful technique for studying magnetic

nanostructures. It is basically an extension of atomic force microscope that images

magnetization patterns with sub-micron resolution [107]. The MFM images show the


44

Figure 3.5 Electron optics of a scanning electron microscope


45

spatial variation of magnetic forces on a sample surface. The changes in the resonant

frequency of the cantilever induced by the magnetic field‟s dependence on the tip-to-

sample separation are detected as shown in the Figure 3.6.

The selection of the tip material is the most important feature for imaging

domains using MFM. High sensitivity to magnetic signal requires a strong magnetic

moment of the tip. However, this high magnetic moment may disturb the actual domain

structure under investigation and consequently the lateral resolution drops with increasing

magnetic moment of the tip. To improve the lateral resolution, thin magnetic coatings are

required. On the other hand, due to decreased magnetic moment of such coatings, the

sensitivity is decreased. An optimization is compromised between lateral resolution and

sensitivity.

In our apparatus, the magnetic tip was coated with cobalt alloy (40 nm thick) on

the tip side and aluminum (30 nm thick) on the detector side, and the tip radius is

typically less than 50 nm. The magnetic tip scans the sample in non-contact mode to

obtain the surface morphology and then a second scan is carried out at constant height

above the surface so that the magnetic and the topographical signals are well separated.

MFM can be used to image naturally occurring and deliberately written domain

structures in the magnetic materials. The images produced by the magnetic tip provide

information about both the topography and the magnetic properties of a surface. The

distance of the tip from the surface determines whether the inter-atomic magnetic force

persists or the Van der Waals force is dominant. As the magnetic force persists for larger

distances, the magnetic effect dominates at larger tip-to-sample distances. If the tip is

close to the surface in the region where standard non-contact AFM is operated, the image

will be predominantly topographic. Magnetic effects become apparent as the separation

between the tip and the sample is increased [108].

3.7 Vibrating Sample Magnetometry

Magnetic characterizations were carried out using a VT-800 (Riken Denshi Co

Ltd.) vibrating sample magnetometer (VSM) with an applied field up to ±15 kOe. In

VSM, the sample is vibrated sinusoidally with frequency „υ‟ and resultantly, a voltage is

induced in suitably placed stationary pickup coils. The output electrical signal of these


46

(a)

(b)

Figure 3.6 (a) Working mode of a schematic MFM with, (b) close image of

magnetic tip and sample surface


47

coils has the same frequency as the source and the intensity is proportional to the

magnetic moment of the sample, the amplitude of vibration, and the frequency [109].

A simplified schematic representation of the VSM is shown in Figure 3.7. The

sample is centered between the poles of a laboratory magnet, capable of generating the

magnetic field Ho. A transducer assembly is located above the magnet and is connected to

the sample holder by a thin vertical rod. The main function of the transducer is to convert

a sinusoidal ac drive signal, provided by an oscillator/amplifier circuit, into a sinusoidal

vertical vibration of the sample rod. This sinusoidal motion is eventually transmitted to

the sample in the uniform magnetic field Ho. The signal resulting from the motion of the

sample is picked up by the coils mounted on the poles of the magnet. The magnitude of

the moment of the sample is proportional to the vibration amplitude and the frequency as

well. Therefore, it is too simplistic to assume that the amplitude is merely proportional to

the magnetization. This difficulty is avoided by employing a nulling technique in which a

vibrating capacitor is used for generating a reference signal that varies with the moment,

vibration amplitude, and vibration frequency in the same manner as the signal from the

pickup coils. An appropriate processing of these two signals makes it possible to

eliminate the effects of vibration amplitude and frequency shifts. In this way, readings are

obtained that vary only with the magnetic moment of the sample. In our case, the M-H

loops were measured with the applied field parallel to the film plane to obtain in-plane

magnetic hysteresis loops.


48

Figure 3.7 Schematic of a typical vibrating sample magnetometer

Pickup

coils

Sample

Coi

ls

Pre-amplifier

Electromagnet

Mechanical

vibrator

Rotation

stage

(Torque) Hall probe

_

+


49

3.8 Superconducting Quantum Interference Device (SQUID)

In recent years, superconducting quantum interference devices (SQUIDs) have

been used in a wide variety of applications due to their superior magnetic field

sensitivity. In our research, a SQUID (MPMS XL, Quantum Design) was employed to

investigate the low temperature magnetic behavior of γ′-Fe4N thin films, and an

externally attached AC resistance bridge (LR-700, Linear Research Inc.) helped us to

investigate the trend of resistivity of the films as the temperature decreases. This external

attachment utilizes “Four Probe Technique” to evaluate the resistivity of the films. A

SQUID system consists of a superconducting ring with a small insulating layer known as

the “weak link” or “Josephson Junction”. The flux passing through the ring is quantized

once the ring has gone superconducting but the weak link enables the flux trapped in the

ring to change by discrete amounts. In this way, the device can be used to measure very

small changes in flux.

The major functional parts of a PC controlled SQUID under MPMS control software

include:

The temperature control module (TCM) which provides an actively regulated,

precision thermal environment over its entire range of operation.

The superconducting magnet system which provides reversible field operation

to ±5 Tesla using an oscillatory technique to minimize magnet drift

immediately following field changes.

The SQUID detector system which includes sensing pick-up loops and

specially designed computer controlled filtering.

The sample handling system which allows automatic sample measurements

and position calibrations using a micro-stepping controller with a positioning

resolution of 0.0003 cm.

The gas handling system which provides gas flow control for temperature

regulation, flushing, and cleaning procedures.

Liquid helium system which provides refrigeration for the superconducting

detection system and magnet, as well as providing for operation down to 1.9 K.


50

Another important part, the sample rod is made with an upper section of needle

temper type 304 stainless steel and a bottom section of quantalloy to minimize the

magnetic signal caused by the sample rod.

3.9 Auger Electron Microscopy (AES)

The estimation of Co, Pt, and Cr at% in iron nitride films was achieved using a

Perkin Elmer-4300 (USA) Auger Electron Spectroscope (AES). A typical AES

spectrometer is schematically is shown in Figure 3.8.

It is usually housed with an ultrahigh vacuum chamber at ≈ 10-10

Torr. This level

of cleanliness is required to prevent surface coverage by contaminants (e.g. C, O) in the

chamber. The electron gun source aims a finely focused beam of (≈ 2 keV) electrons at

the specimen surface, where it is scanned over the region of interest. Energy of the

emitted Auger electrons is analyzed by a cylindrical analyzer which consists of coaxial

metal cylinders raised to different potentials. The electron pass energy (E) is proportional

to the voltage on the outer cylinder while the incremental energy range (∆E) of

transmitted electrons determines the resolution (∆E/E), which is typically 0.2 % to 0.5 %

[110].

Electrons with higher or lower energies than E hit either the outer or inner

cylinders, respectively. They do not exit the analyzer and are not counted. By sweeping

the bias potential on the analyzer, the entire electron spectrum is obtained. Auger

electrons are but a part of the total electron yield, N(E), intermediate between low energy

secondary and high energy elastically scattered electrons. They are barely discernible as

small bumps above the background signal. Therefore, to accentuate the energy and

magnitude of the Auger peaks, the spectrum is electronically differentiated and this gives

rise to the common AES spectrum, or dN(E)/dE vs E response [111].

The quantitative at% estimation was evaluated from Survey Scan Data obtained

from AES spectra. Atomic concentration Cx of an element „x‟ may be estimated by

)2.3......(........................................ SI

SIC xx

x


51

Figure 3.8 Schematic of auger electron spectroscope


52

where, Ix is the peak to peak height of the Auger electron signal for the element „x‟

divided by the proper scale sensitivity and Sx is the sensitivity factor. The AES spectra

were obtained at a base pressure of 8.0 × 10-10

Torr and a working gas (Ar: for pre-

sputtering) pressure of 9.0 × 10-8

Torr. Each sample was pre-sputtered for 24 seconds to

remove any contamination and oxide layers upto few nanometers from the sample

surface.

3.10 Transmission Electron Microscope

3.10.1 TEM-Microstructure

Micro structural investigations, Lorentz microscopy and electron holography were

performed by a Tecnai G2 F30 S-Twin, transmission electron microscope (TEM). The

TEM has been considered an indispensible tool for characterizing microstructure, so

several TEM techniques have been developed to study magnetic materials. The schematic

illustration of a TEM system is presented in Figure 3.9.

The electrons are provided by an electron gun. The filament consists of tungsten

or LaB6 crystals, which is accordingly heated and emits electrons. The electrons in a field

emission gun (FEG), however, are extracted from an extremely pointed cathode by an

electric field. The sharp tip causes a strong electric field at its very end, as the field is

inversely proportional to the radius of the tip‟s curvature. Therefore, electrons are emitted

from a very limited area which in turn provides an excellent spatial coherence of the

electron wave, a crucial requirement for electron holography. The emitted electrons are

accelerated in a multistage process to reach their final highly relativistic energy of several

hundred keV, depending on the microscope. A higher acceleration voltage causes a

smaller de-Broglie wavelength λ of the electrons, typically in the pm range. The high

energy of electrons ensures sufficient penetration power to penetrate thin metallic films

with typical thickness of ≈ 100 nm. The specimen is inserted into the electron microscope

with a specimen holder, which restricts the dimensions of the sample to approximately

2mm [112].

A fluorescent screen allows the direct observation of the image. It can be recorded

with a conventional plate camera, although most of the microscopes are equipped with a

CCD camera, allowing digital acquisition of the image.


53

Figure 3.9 Typical components of a Transmission Electron Microscope, equipped with

a Lorentz lens


54

3.10.2 Lorentz Microscopy

In a common set up of a microscope, the sample sits right inside the objective

lens, which often acts as a combination of a further condenser lens and an imaging lens.

The magnetic field in the fully excited objective lens exceeds 1 T and would severely

influence the micromagnetic configuration of the specimen. Therefore, the objective

cannot be used for magnetic imaging techniques. However, a low magnification mode is

usually implemented in all microscopes, which operates with the objective lens turned

off. Its action is performed by another lens called diffraction lens, which is located at a

larger distance from the specimen. This requires a long focal length lens, which yields

only a very moderate magnification but at the same time, causes only a negligible

magnetic field at the specimen‟s location.

Another important approach is to use a special Lorentz lens, which serves as an

objective lens, as shown in Figure 3.9. This Lorentz lens is located at a larger distance

from the sample than the objective lens but still closer than the lenses usually in use. The

Lorentz lens allows investigations with virtually zero magnetic field in the sample‟s

region. This set up allows magnification in the 105 range and the resolution can reach up

to 2 nm. Several lenses follow which magnify the Lorentz TEM image further. Lorentz

microscopy has the advantage to exclude in-elastically scattered electrons from the image

which in turn allows reducing the background noise and opening the opportunity to

investigate even thicker specimens or samples with weak magnetic inductions.

3.10.3 Electron Holography

Electron holography is an alternative electron microscopy technique that has

proved ideal for studying the magnetic state and response of ferromagnetic materials at

length scales approaching the nanometer level. By measuring the phase change of the

electron wave that has traveled through the sample relative to an unperturbed reference

wave, a quantitative measure of the local magnetic field can be obtained. This technique

relies on the interference between two coherent electron waves, as shown schematically

in Figure 3.10. The FEG provides the essential beam coherence and the electrostatic

biprism, normally located in the selected-area aperture plane, is used to interfere the wave

scattered by the object with the reference vacuum wave. The resulting interference


55

Figure 3.10 Schematic ray diagram illustrating experimental geometry used for

electron holography in TEM


56

pattern is then processed in order to retrieve the electron phase, in turn allowing access to

desired information about the magnetic field of the sample. This technique can in

principle achieve a spatial resolution of better than 1 nm for magnetic materials [113].

3.11 Magneto-Optic Microscope Magnetometery

The MOMM system, as shown in Figure 3.11, is a magnetometric setup and

basically an extension of magneto-optical Kerr microscope, which is a well established

technique extensively used in magnetic domain observations [114].

The MOMM basically consists of a polarizing optical microscope set to visualize

in-plane magnetic domain images via magnetic contrast, utilizing a longitudinal magneto-

optical Kerr effect (MOKE). The optical illumination path is tilted for the longitudinal

MOKE to provide an incident angle of 20˚ from the film normal by shifting the position

of the objective lens as well as adjusting the relevant optics. The Kerr intensity signal is

detected by a digital intensified charge-couple device (CCD) camera system with a

maximum gain of 106. The MOMM system has a spatial resolution of 400 nm and the

Kerr angle resolution is 0.1˚. An electromagnet, positioned just below the sample stage, is

controlled by a PC to apply an external magnetic field to the sample in the range of

±3.5 kOe.

The magnetic domain images are stored with the help of a digital video

processing board having an image grabbing rate of 30 frames/sec in real time. The time

resolved domain images on a 400×320 μm2 sample area are initially grabbed on a 256

grey scale, and then intensified by advanced image processing techniques, such as

background subtraction, noise filtering, and black-and-white image extraction.

This system has the ability to probe into the local dynamic behavior during

magnetization reversal and to characterize statistically the Barkhausen effect [152] which

states that the reversal of magnetization takes place in form of discrete and jerky jumps

rather than to proceed smoothly. The visualization capability of the MOMM enables us to

directly visualize and investigate the motion of domain wall in the Barkhausen

avalanches. The repeated observation of the domain-wall motion reveals that there exist

some pinning segments around which domain walls are very flexible. The flexible part of

the domain wall moves forward via a Barhausen jump, while the pinned part is fixed at


57

Figure 3.11 Schematic of a Magneto-Optical Microscope magnetometer


58

the same position for a relatively long time.

The Barkhausen avalanche is triggered by applying a constant magnetic field to

an initially saturated sample. The strength of applied field is constant near the coercive

field to eliminate the influence caused by the difference in the field-sweeping rates. By

means of this setup, the Barkhausen jumps are directly visualized and characterized from

serial, time resolved domain images.

Chapter 4 Preparation and Characterization of γ΄-Fe4N Thin Films

59

Preparation and Characterization of γ′-Fe4N Thin

Films

(Results and Discussions I)

γ′-phase of iron nitride is one of the most stable phases among Fe-N family and

bears a simple face-centered cubic (fcc) structure with nitrogen atom positioned at the

body-centered site. This phase has a high saturation magnetic moment and is an attractive

choice for possible use in magnetic devices. Therefore, the interest in γ′-Fe4N thin films

has increased manifolds during recent years. Some attempts have already been made to

deposit single phase epitaxial films of this phase in order to exploit this material for its

potential applications in high density magnetic write heads and magnetic recording

media.

Furthermore, it is expected to replace those materials for which high magnetic

flux density and low coercivity is required, e.g. 50% nickel permalloys [50]. In case of

current perpendicular-to-plane (CPP) devices, it could also be an attractive material,

where low resistance of the ferromagnetic materials could be an issue and the possible

development of all nitride all epitaxial magnetic tunnel junctions (MTJs) with epitaxial

Cu3N as an insulating layer [115] between epitaxial γ΄-Fe4N magnetic electrodes.

Recently developed MgO-based MTJs with γ΄-Fe4N electrode have been reported having

relatively large inverse tunneling magnetoresistance (TMR) effect [116]. This

development has further triggered interest in γ΄-Fe4N films.


60

4.1 Preparation of γ′-Fe4N Thin Films

In order to find the optimum conditions for single phase epitaxial growth of γ′-

Fe4N thin films, several iron nitride thin films were grown on MgO(100) substrates using

reactive dc magnetron sputtering with varying deposition parameters. The base pressure

of the chamber was better than 2 x 10-6

Torr. While depositing different series of

samples, nitrogen partial pressure was varied from 0.2 to 2.0 mTorr, substrate

temperature was varied from 200 to 550 ˚C, annealing time was changed from 10 minutes

to 2 hours, dc sputtering power was changed from 15 to 60 W, and film thickness was

varied from 20 to 200 nm. After several attempts, employing different combinations of

parameters, following series were found to approach the favorable conditions for single

phase growth. Therefore, these series were re-deposited.

(a) Changing nitrogen partial pressure from 0.2 to 1.0 mTorr at 400 ˚C substrate

temperature and in-situ annealing for 30 minutes. The films were deposited at a

deposition rate of 0.55 Å/sec utilizing dc sputtering power of 30 W. The thickness

of each film was 50 nm. Complete deposition conditions are shown in Table 4.1

(b) Changing substrate temperature from 200 to 500 ˚C and in-situ annealing for

30 minutes, with fixed nitrogen partial pressure at 0.5 mTorr. The films were

deposited at a deposition rate of 0.55 Å/sec utilizing dc sputtering power of 30 W.

The thickness of each film was 50 nm. Complete deposition conditions are shown

in Table 4.2

(c) Changing annealing time from 10 to 40 minutes with fixed substrate temperature at

450 ˚C, and nitrogen partial pressure at 0.5 mTorr. The films were deposited at a

deposition rate of 0.55 Å/sec utilizing dc sputtering power of 30 W. The thickness

of each film was 55 nm. Complete deposition conditions are shown in Table 4.3

(d) Changing substrate material, with fixed substrate temperature at 450 ˚C, annealing

time at 30 minutes, and nitrogen partial pressure at 0.5 mTorr. The films were

deposited at a deposition rate of 0.55 Å/sec utilizing dc sputtering power of 30 W.

The thickness of each film was 55 nm.


61

Table 4.1 Nitrogen partial pressure dependent preparation of γ′-Fe4N films, Series 1

DC Sputtering Power = 30 W

PAr = 5.0 mT

Film Thickness= 50 nm

Deposition Time = 15 min.

Sample No 2NP ( mTorr) Substrate Temp.

(˚C)

Annealing Time

(Minutes)

Substrate Material

1(a) 0.0 400 30 MgO(100)

1(b) 0.2 400 30 MgO(100)

1(c) 0.4 400 30 MgO(100)

1(d) 0.6 400 30 MgO(100)

1(e) 0.8 400 30 MgO(100)

1(f) 1.0 400 30 MgO(100)


62

Table 4.2 Substrate temperature dependent preparation of γ′-Fe4N films, Series 2


PAr = 5.0 mT




(˚C)

Annealing Time

(Minutes)

Substrate Material

2(a) 0.5 200 30 MgO(100)

2(b) 0.5 250 30 MgO(100)

2(c) 0.5 300 30 MgO(100)

2(d) 0.5 350 30 MgO(100)

2(e) 0.5 400 30 MgO(100)

2(f) 0.5 450 30 MgO(100)

2(g) 0.5 500 30 MgO(100)


63

Table 4.3 Annealing time dependent preparation of γ′-Fe4N films, Series 3


PAr = 5.0 mT




(˚C)

Annealing Time

(Minutes)

Substrate Material

3(a) 0.5 450 10 MgO(100)

3(b) 0.5 450 20 MgO(100)

3(c) 0.5 450 30 MgO(100)

3(d) 0.5 450 40 MgO(100)


64

4.2 Crystal Phase Determinations

4.2.1 X-ray Diffraction Analysis

After sample preparation, the next and the most important step is crystal phase

determination. For this purpose, XRD is the most widely used technique. Owing to the

huge data bank available covering practically every phase of every known material

(powder diffraction patterns), it is routinely possible to identify phases in crystalline bulk

and thin film (≥ 100 Å) materials.

In our work, the investigation of crystal structure was carried out using Rigaku

D/MAX-RC MPA x-ray diffractmeter (XRD) under the operating conditions as explained

in section 3.3.1, and the data was analyzed using MDI/JADE5 software.

The effect of nitrogen (obtained from ammonia source) on the crystal structure of

iron nitride films have been described by Yamaguchi et al. [33]. We have used nitrogen

obtained from pure source in order to investigate the crystalline quality of the iron nitride

films.

Figure 4.1 shows the XRD patterns of thin films deposited on MgO(100)

substrates with varying nitrogen partial pressure (2NP ). In this series, the substrate

temperature was 400 ˚C and all the films were in-situ annealed for 30 minutes. The XRD

patterns revealed that nitrogen partial pressure strongly affected the iron nitride phase

formation. Figure 4.1 (b) shows the XRD pattern of pure Fe deposited on MgO(100)

substrate without nitrogen flow, therefore the unique peak of α-Fe (200) at 2θ = 65.06°

was observed. All other peaks were from MgO(100) substrate. For comparison, the XRD

pattern of single crystal MgO(100) substrate is provided in Figure 4.1 (a). A low intensity

peak of γ΄-Fe4N (200) along with α-Fe (200) at 2θ = 65.06° appeared at 2NP = 0.2 mTorr

as shown in Figure 4.1 (c). As the 2NP was increased to 0.4 mTorr, the intensity of γ΄-

Fe4N (200) peak increased and an additional peak of γ΄-Fe4N (100) at 2θ = 23.45°

appeared as well as shown in Figure 4.1 (d). The corresponding intensity of α-Fe (200)

peak decreased in this sample. This reveals that with the injection of more nitrogen, more

iron atoms have reacted with nitrogen and accordingly have converted into γ΄-Fe4N.


65

Perfect match between nitrogen and iron atoms was observed at 2NP = 0.6 mTorr when no

iron peak was seen and only (100) and (200) peaks of γ΄-Fe4N were observed at their

perfect 2θ values revealing a high level of single phase epitaxial growth (Figure 4.1e).

Further increase in nitrogen partial pressure up to 0.8 – 1.0 mTorr, could not favor the

epitaxial growth of γ΄-Fe4N films as some small unknown nitride peaks were observed

along with a small iron peak appearing again as can be seen in the XRD pattern in Figure

4.1 (f-g). This means that 0.6 mTorr is the optimum partial pressure of nitrogen for single

phase epitaxial growth of γ΄-Fe4N films on MgO(100) substrates at 400 ˚C with in-situ

annealing of 30 minutes.


substrates with varying substrate temperature from 200 – 500 ˚C. The nitrogen partial

pressure (2NP ) was kept at 0.5 mTorr and all the films were in-situ annealed for 30

minutes. The XRD patterns revealed that the substrate temperature affected the iron

nitride phase formation, as well. Figure 4.2 (b) shows the XRD pattern of the sample

deposited on MgO(100) at a substrate temperature of 200 ˚C. The pattern showed only a

small intensity peak of γ΄-Fe4N (200) at 2θ = 47.91°. All the other peaks were from

MgO(100) substrate as can be compared with the XRD pattern of the pure single crystal

substrate provided in Figure 4.2 (a). As the substrate temperature was increased to 250˚C,

the intensity of the (200) peak of γ΄-Fe4N increased and a small intensity (100) peak of

γ΄-Fe4N also appeared at 2θ = 23.45° as shown in Figure 4.2 (c). The intensities of both

the (100) and (200) peaks of γ΄-Fe4N were increased by the further increase of

temperature (with a step of 50 ˚C) and became maximum at a temperature of 450 ˚C as

can be seen in the Figures 4.2 (d-g). When the substrate temperature was further

increased to 500 ˚C, the intensity of both the (100) and (200) peaks of γ΄-Fe4N were

decreased along with the appearance of (200) peak of α-Fe at 2θ = 65.06° (Figure 4.2f). It

is evident that the increase of substrate temperature up to 450 ˚C has favored the

formation of γ΄-Fe4N phase to a high quality single phase epitaxial texture, while further

increase in temperature has resulted in the collapse of epitaxial phase due to appearance

of an α-Fe peak. This probably is due to the fact that a high temperature has resulted in

excessive escape of nitrogen from the film causing the appearance of the α-Fe peak.


66

Figure 4.1 XRD patterns of (a) Pure MgO(100) substrate, (b) Pure α-Fe deposited on

MgO(100), Iron Nitride samples deposited using nitrogen partial pressure

of (c) 0.2 mTorr, (d) 0.4 mTorr, (e) 0.6 mTorr, (f) 0.8 mTorr,

and (g) 1.0 mTorr

20 30 40 50 60 70

(a)

2()

(b)

(c)

(d)

(e)

(f)

-Fe4N (100)

-Fe4N (200)

-Fe (200)

Unknown Nitride

(g)


67

Figure 4.2 XRD patterns of (a) Pure MgO(100) substrate, γ΄-Fe4N thin films deposited

on MgO(100) substrates at deposition temperatures of (b) 200 ˚C, (c) 250 ˚C,

(d) 300 ˚C, (e) 350 ˚C, (f) 400 ˚C, (g) 450 ˚C, and (h) 500 ˚C

20 30 40 50 60 70

2()

(a)

(b)

(c)

(d)

-Fe4N (100)

-Fe4N (200)

-Fe (200)

(e)

(f)

(g)

(h)


68


substrates with annealing time varying from 10 – 40 minutes [50]. The nitrogen partial

pressure (2NP ) was kept at 0.5 mTorr and all the films were deposited at the substrate

temperature of 450 ˚C. The XRD patterns revealed that annealing time has conclusively

affected the iron nitride phase formation. Figure 4.3 (a) shows the XRD pattern of the

sample deposited on MgO(100) at a substrate temperature of 450 ˚C and in-situ annealed

for 10 minutes. The XRD pattern clearly illustrated (100) and (200) peaks of γ΄-Fe4N at

2θ of 23.45° and 47.91°, respectively. All other peaks were from MgO substrate. For

comparison, the x-ray diffraction pattern of pure single crystal MgO substrate used in this

work is also shown in Figure 3(e). The intensity of (100) and (200) peaks of γ΄-Fe4N

films increased as the annealing time was increased to 20 minutes. A perfect single phase

epitaxial texture was observed at an annealing time of 30 minutes as the maximum

intensities of (100) and (200) peaks of γ΄-Fe4N film were observed under these

conditions. This epitaxial single phase texture was no more there when the annealing time

was increased to 40 minutes as some extra peaks of α-Fe (200) at 2θ = 65.06° and α˝-

Fe16N2 (220) at 2θ = 44.77° were also observed as shown in XRD pattern in Figure 3(d).

This result might be attributed to the excessive escape of nitrogen atoms from the iron

nitride texture due to excessive annealing that eventually caused the conversion of some

of the molecules of 4 : 1 iron nitride phase, partially into 8 : 1 iron nitride phase and

partially into pure iron.

After studying the XRD patterns of these three series of samples, it was evident

that nitrogen partial pressure of 0.5 mTorr, substrate temperature of 450 °C and annealing

time of 30 minutes are the most suitable conditions for epitaxial single phase growth of

γ΄-Fe4N films


69

.

Figure 4.3 XRD patterns of γ΄-Fe4N thin films deposited on MgO(100) substrates at

deposited at 450 ˚C and in-situ annealed (a) 10 min. (b) 20 min. (c) 30 min.

(d) 40 min. and (e) XRD pattern of pure MgO(100) substrate

20 30 40 50 60 70

2

(e)

(d)

-Fe4N (100)

-Fe4N (200)

-Fe (200)

-Fe16N2 (220)

(c)

(b)

(a)


70

4.2.2 Rocking Curve Analysis

The rocking curve (RC) analysis is a very useful technique employed for the

observation of the preferred orientations in the epitaxially gown films. The value of full-

width at half maximum (FWHM) of the RC is used as a measure of degree of the

preferred orientations.

In our study, this analysis was carried out to investigate the effect of annealing

time on the degree of epitaxy of γ′-Fe4N films. Figures 4.4 – 4.7 show the rocking curves

obtained for the (100) reflection of the samples annealed for 10 to 40 minutes. Figure 4.8

shows graphically the effect of annealing time on the FWHM value of the samples

annealed for 10-40 minutes. It is obvious that the value of FWHM decreased as the

sample is annealed from 10 to 30 minutes but suddenly rose for the sample annealed for

40 minutes. The RC measurements were carried out performing ω-scan at a fixed 2θ =

23.45˚. The rocking curve obtained for the as deposited sample demonstrated the full

width at half maximum value of 0.74˚. As the sample was annealed for 10 minutes, the

peak intensity was increased and the value of FWHM decreased to 0.66˚. The trend

continued up to 30 minutes annealing for which minimum FWHM value of 0.37˚ and the

maximum peak intensity was observed. This value of FWHM is better than reported by

Tamura et al. [117] for this material. While, this trend was suddenly changed for the

sample annealed for 40 minutes as the peak intensity was decreased and the FWHM

value increased to 1.46˚. Therefore, these results demonstrate that 30 minutes annealing

at 450 °C substrate temperature are the most favorable conditions for single phase

epitaxial growth of γ΄-Fe4N thin films on MgO substrate.


71

Figure 4.4 Rocking curve of γ′-Fe4N films obtained for (100) reflection for the sample

annealed for 10 minutes


annealed for 20 minutes.

18 20 22 24 26 28

0

500

1000

1500

2000

Inte

nsi

ty (

a.u

)

Angle (°)

FWHM = 0.6606

(a)

18 20 22 24 26 28

0

1000

2000

3000

4000

5000

Angle (°)

Inte

nsi

ty (

a.u

.)

FWHM = 0.52

(b)


72





18 20 22 24 26 28

0

200

400

600

800

1000

1200

Angle (°)

Inte

nsi

ty (

a.u

.)

FWHM = 1.46

(d)

18 20 22 24 26 28

0

1000

2000

3000

4000

5000

6000

Angle (°)

Inte

nsi

ty (

a.u

.)

FWHM = 0.37

(c)


73

Figure 4.8 FWHM plotted against annealing time of γ′-Fe4N thin films

0 10 20 30 40

0.3

0.6

0.9

1.2

1.5

FW

HM

(

)

Annealing Time (min.)


74

4.2.3 Phi Scan Analysis

The phi (ϕ) scan provides us information on the in-plane epitaxy because any set

of lattice planes perpendicular to the surface can be brought into diffraction. In our study,

the in-plane epitaxy of the γ′-Fe4N sample, annealed for 30 minutes was further verified

by the phi scan obtained using x-ray diffractometry and is shown in Figure 4.9. The scan

was performed at χ = 54.7º (which is angle between the (100) and (111) planes in a cubic

system) from 0º to 360º [91], in steps of 1º, at a fixed 2θ and ω. The scan was performed

for (200) reflection of γ′-Fe4N with the fixed values of 2θ at 47.91º and ω at 23.96º. The

appearance of only four sharp peaks at exact distance of 90˚ apart in Φ-scan with 2θ fixed

at (111) reflection evidently prove the high quality epitaxial nature of γ΄-Fe4N thin films.

4.2.4 Pole Figure Analysis

The in-plane orientation of the film on MgO(100) substrate was also investigated

and confirmed by the pole figure measurements as shown in Figure 4.10. The sample was

rotated around the ϕ axis from 0º to 360º in steps of 1º, again at fixed 2θ = 47.91º and

ω = 23.96º. The presence of four clear spots at exact angular displacement of 90˚ in the x-

ray ψ-Φ pole scan with 2θ fixed at (111) reflection confirmed the high quality in-plane

epitaxial texture of γ΄-Fe4N thin films, as was evidenced by the phi-scan analysis.


75

0 45 90 135 180 225 270 315 3600

50

100

150

200

Inte

nsity (

a.u

.)

Angle

Figure 4.9 The phi (ϕ) scan for (200) reflection of γ′-Fe4N films obtained for the sample



76

Figure 4.10 The pole figure for (200) reflection of γ′-Fe4N films obtained for the sample



77

4.3 Surface Morphology

In order to investigate the surface morphology of films, the scanning electron

microscope (SEM) provides researchers with a highly magnified image of the thin film

surface that is very similar to what one would expect to actually "see" the surface

visually. The features of surface morphology like film surface topography and

microstructure in plain view including grain size and shape, existence of compounds,

presence of hillocks or whiskers, or evidence of film voids, are important in determining

the film quality.

A regular nano-sized grain structure was imaged with the help of scanning

electron microscope for the γ′-Fe4N films deposited at 200 ºC substrate temperature and

in-situ annealed for 30 minutes as shown in Figure 4.11. A relative increased grain size

was observed for the sample deposited at 300 ºC substrate temperature (Figure 4.12).

This increment in grain size is attributed to the increased surface mobility of the adsorbed

species caused by the relative higher substrate temperature. High surface mobility is

included among one of the factors required for epitaxial growth of thin films [118].

Further increase in substrate temperature up to 400 ºC resulted in the coalescence of the

larger grains into the tunnel like structures as evidenced by the SEM micrograph shown

in Figure 4.13. Finally, at 450 ºC substrate temperature, a continuously grown smooth

surface image was observed (Figure 4.14) for the γ′-Fe4N film which was a direct

evidence of layer by layer growth mechanism. Quite often, the surface mobility is

reduced at critically higher temperature as was witnessed in our case as well, so a reduced

sized grains were again visible for the sample deposited at 500 ºC substrate temperature

as shown in Figure 4.15.

In this study, the determination of accurate volume of the iron nitride films is very

crucial as it is used for estimation of volumetric saturation magnetization. Therefore, the

thickness calibration performed initially with the help of atomic force microscope was

finally confirmed by cross-sectional images captured using scanning electron microscope

as shown in the Figures 4.16 – 4.19


78

Figure 4.11 Scanning Electron Microscope image of γ′-Fe4N sample deposited at 200 ˚C

substrate temperature




79






80




81

Figure 4.16 Cross-sectional SEM image of γ′-Fe4N sample in-situ annealed for 10 min



82


Figure 4.19 Cross-sectional SEM image of γ′-Fe4N sample in-situ annealed for 40 min.


83

4.4 RMS Surface Roughness

In addition to thickness calibration, atomic force microscope was also employed

to investigate the topography and the subsequent surface roughness of the γ′-Fe4N films.

The microstructure and topographical details of a thin film of any material depend on the

kinetics of growth and hence mainly on the substrate temperature and post deposition

annealing time. A smooth surface film might favor epitaxial growth. For smooth layer by

layer growth of the films, high surface mobility is required which can be achieved in two

ways: high substrate temperature and the thermal annealing of the films [118].

In this study, the minimum root mean square (RMS) value of surface roughness

was achieved by optimizing the annealing time for the series of samples deposited at 450

˚C substrate temperature. Figures 4.20 – 4.23 show the topographical images of the γ′-

Fe4N films deposited on MgO(100) and in-situ annealed for 10 – 40 minutes. Figure 4.24

shows the effect of annealing time on the RMS value of surface roughness. The value of

surface roughness for the as-deposited sample was observed to be 0.4 nm. The RMS

roughness decreased to 0.3 nm when the sample was in-situ annealed for 10 minutes.

This decrease in roughness is attributed to the increased surface mobility of deposited

species caused by the increased annealing time. The trend of decrease in surface

roughness continued up to the sample annealed for 30 minutes which exhibited minimum

RMS value of 0.17 nm. Such a low value of RMS roughness clearly demonstrates

epitaxial growth from smooth layer-by-layer growth process, in which, first islands are

formed which then coalescence into full smooth layered texture [44].

The sample annealed for 40 minutes demonstrated a drastic change in structure:

the values of RMS roughness increased suddenly to 0.7 nm. This abrupt change in

behavior of the film annealed for 40 minutes is obviously due to the non-epitaxial

growth, which is consistent with the XRD result of the sample as can be seen in Figure

4.3 (d). This behavior is attributed to the fact that longer annealing time has resulted in

more than sufficient increase in surface mobility which in turn has interrupted the layer

by layer growth process and hence favoring non-epitaxial growth.


84

An interesting relationship was evident when we plotted the effect of annealing

time against the RMS value of surface roughness and the value of FWHM using double

Y axis as shown in Figure 4.25. The RMS value of surface roughness and FWHM

showed exactly the same behavior and decreased with the increase of annealing time up

to 30 minutes. The lowest value of surface roughness (0.17 nm) and minimum FWHM

value of 0.37° were observed for this sample. The trend changed drastically for the

sample annealed for 40 minutes for which the values of RMS roughness and the FWHM

increased suddenly to 0.7 nm and 1.46°, respectively. Therefore, a prominent effect of

annealing time was observed on the RMS roughness and FWHM values of the samples.

From these observations it can clearly be deduced that 450 ˚C substrate temperature

followed by in-situ annealing of 30 minutes are the most favorable conditions for single

phase epitaxial growth of γ′-Fe4N thin films.


85

Figure 4.20 Topographical image obtained using AFM for the sample annealed for

10 min


20 min


86


30 min


40 min


87

Figure 4.24 RMS roughness plotted against annealing time of γ′-Fe4N thin films

0 10 20 30 40

0.2

0.4

0.6

0.8

RM

S R

oughne

ss (

nm

)



88

0 10 20 30 40

0.0

0.2

0.4

0.6

0.8


RM

S R

oughne

ss (

nm

)

0.0

0.5

1.0

1.5

2.0

2.5

FW

HM

(

)

RMS Roughness (nm)

FWHM ()

Figure 4.25 RMS roughness and FWHM plotted against annealing time

of γ′-Fe4N thin films


89

4.5 Magnetic Characterization

4.5.1 Saturation Magnetization

Vibrating Sample Magnetometery is one of the most popular and reliable

technique used for magnetic characterization. We employed this technique for the

measurement of saturation magnetization (Ms) of γ′-Fe4N thin films deposited on

MgO(100) at a substrate temperature of 450 ˚C and in-situ annealed up to 40 minutes.

Figure 4.26 shows M-H loops achieved with the help of vibrating sample

magnetometer for the samples in-situ annealed for 10 to 40 minutes. Figure 4.27 shows

the plot which reveals the effect of annealing time plotted against Ms and RMS value of

surface roughness for γ΄-Fe4N thin films deposited on MgO(100) substrate at 450 °C

temperature and in-situ annealed up to 40 minutes. The value of saturation magnetization

for the as-deposited sample was 1510±20 emu/cc with a maximum value of

corresponding surface roughness. As the sample was annealed for 10 minutes, the Ms

value increased to 1590±20 emu/cc. The trend continued and the value of saturation

magnetization increased with the decrease in RMS roughness and increase in annealing

time up to the sample annealed for 30 minutes showing a value of 1760±20 emu/cc. This

value is 16% higher than the previous reported value [119] and might be attributed to the

high degree of epitaxy for this sample. The sample annealed for 40 minutes showed

sudden decrease in Ms value (1140 emu/cc) which is obviously due to the non-epitaxial

nature and high RMS value of surface roughness (0.7 nm) of the sample. Table 4.4 shows

the relationship between annealing time and the corresponding values of RMS roughness

(nm), FWHM (θ), Ms (emu/cc), and μB/Fe.

4.5.2 FWHM Dependent Ms

Figure 4.30 (a) shows a plot of Ms vs FWHM for γ΄-Fe4N thin films deposited on

MgO(100) substrates at a temperature of 450 °C and in-situ annealed for 10 to

40 minutes. As seen in the figure, the value of saturation magnetization increases with the

decrease in FWHM up to the sample annealed for 30 minutes showing a value of

1760±20 emu/cc. As stated in section 4.5.1, this value is 16% higher than the previous


90

-10000 -5000 0 5000 10000-2

0

2

me

mu

H (Oe)

(a)

-10000 -5000 0 5000 10000-2

0

2

me

mu

H (Oe)

(b)

-10000 -5000 0 5000 10000

-2

0

2

me

mu

H (Oe)

(c)

-10000 -5000 0 5000 10000

-1

0

1

me

mu

H (Oe)

(d)

Ms vs H at 10, 20, 30 & 40 min annealing

Figure 4.26 M-H loops obtained using VSM for the samples annealed for (a) 10 min.

(b) 20 min. (c) 30 min. and (d) 40 min.


91

0 10 20 30 40

0.2

0.4

0.6

0.8

Annealing Time (min)

RM

S R

ou

gh

ne

ss

(nm

)

1200

1400

1600

1800

Ms

(em

u/c

c)

Figure 4.27 Effect of annealing time plotted against RMS roughness and saturation

magnetization of the γ΄-Fe4N films


92

Table 4.4 A table showing annealing time and the corresponding values of RMS

roughness, FWHM, Ms, and μB/Fe

Sr. No. Annealing

Time (min.)

RMS roughness

(nm)

FWHM (θ) Ms (emu/cc) μB/Fe

1. 0 0.4 0.74 1510 2.22

2. 10 0.3 0.66 1590 2.34

3. 20 0.25 0.52 1680 2.47

4. 30 0.17 0.37 1760 2.59

5. 40 0.7 1.46 1140 1.68


93

reported value [119] and might be attributed to the high degree of epitaxy for this sample.

The sample annealed for 40 minutes shows sudden decrease in Ms value (1140 emu/cc),

which is obviously due to the non-epitaxial nature of the film.

4.6 Substrate Dependent Preparation of γ΄-Fe4N

4.6.1 Preparation

To further explore the dependence of saturation magnetization on epiatxial

growth, we have deposited γ΄-Fe4N thin films under the same deposition conditions on

three different substrates: MgO(100), SrTiO3(100), and LaAlO3(100) with lattice

mismatches of 11, 3, and 0%, respectively. All the films were deposited under the same

conditions: at a nitrogen partial pressure of 0.5 mTorr, substrate temperature of 450 ˚C,

and in-situ annealing of 30 minutes. Detailed preparation conditions are shown in Table

4.5.

4.6.2 XRD Analysis

Figures 4.28 (b), (d), and (f) show the XRD patterns of the films deposited on

MgO(100), SrTiO3(100), and LaAlO3(100) substrates, respectively; while the Figures

4.28 (a), (c), and (e) show the three corresponding XRD patterns of pure single crystal

substrates for comparison. Although MgO has 11% lattice mismatch with γ΄-Fe4N, even

then favorably can be employed for epitaxial growth as two clear (100) and (200) peaks

of γ΄-Fe4N were evident as revealed by the XRD pattern shown in Figure 4.28 (b), but

such a large mismatch could affect the growth normals and the in-plane orientation

relationship between the substrate and the deposit [42].

Having lattice mismatch of only 3%, SrTiO3 offers relatively better template for

epitaxial growth. X-ray diffraction pattern of the sample deposited on SrTiO3(100)

showed unique (200) peak of γ΄-Fe4N at 2θ value of 47.91° (Figure 4.28d), indicating the

epitaxial structural texture of the film. For perfectly matched lattice, LaAlO3 should be

the most appropriate substrate for the growth of epitaxial γ΄-Fe4N thin films. One

problem however was that, due to perfect lattice matching, the dominant peaks of both

the deposit and the substrate completely overlapped each other hence could not be


94

Table 4.5 Substrate dependent preparation of γ′-Fe4N films, Series 4


PAr = 5.0 mT




(˚C)

Annealing Time

(Minutes)

Substrate Material

4(a) 0.5 450 30 MgO(100)

4(b) 0.5 450 30 SrTiO3(100)

4(c) 0.5 450 30 LaAlO3(100)


95

Figure 4.28 XRD patterns of (a) MgO(100) substrate, (b) γ΄-Fe4N on MgO,

(c) SrTiO3(100) substrate, (d) γ΄-Fe4N on SrTiO3, (e) LaAlO3(100)

substrate, and (f) γ΄-Fe4N on LaAlO3

20 30 40 50 60 70

-Fe4N (100)

-Fe4N (200)

-Fe (200)

2

MgO

SrTiO3

(f)

(e)

(d)

(c)

(b)

(a)

LaAlO3


96

detected by XRD patterns in case of thin films as shown in Figure 4.28 (f). At the same

time it is inferred that perfect epitaxial single phase texture of γ΄-Fe4N thin films might

have been deposited on LaAlO3(100) substrate, as no peak of any other phase was

obvious on the XRD pattern.

4.6.3 Substrate Dependent Ms

Single phase thin films of α″-Fe16N2 have been known to manifest the highest

values of saturation magnetization as claimed by Kim and Takahashi [12], and then

confirmed by Sugita and co-workers [53-58] and Huang et al. [63]. All these thin films

were epitaxially grown, so the higher values of saturation magnetization were attributed

to the epitaxial growth. In our research, we have tried to explore the epitaxy dependent

saturation magnetization of γ′-Fe4N films. For the growth of epitaxial films, one of the

most important conditions is the crystallographic compatibility between the substrate and

the deposit. Therefore, single crystal substrates with minimum lattice mismatches are

always an appropriate choice for epitaxial growth.

Figure 4.29 shows the M-H loops of γ′-Fe4N films deposited on three different

single crystal substrates: MgO(100), SrTiO3(100), and LaAlO3(100) having 11%, 3%,

and 0% lattice mismatches with γ′-Fe4N, respectively.

The values of saturation magnetization observed for these three samples are

plotted as function of lattice mismatch as shown in Figure 4.30 (b). It can be shown that

saturation magnetization increases with decrease of lattice mismatch and reached up to

1980±20 emu/cc at 0% lattice mismatch for LaAlO3 substrate. Note that this value of

saturation magnetization is about 24% higher than the experimentally observed value

[119] reported so far for this material. These results clearly demonstrate that lattice

mismatch has a strong impact on the epitaxial growth of γ′-Fe4N films which in turn has

prominent influence on the saturation magnetization of iron nitride films.

Table 4.6 shows the lattice constants of the deposit, the substrates, lattice

mismatches, the Ms, and μB/Fe values of γ′-Fe4N films.


97

Figure 4.29 M-H loops for the γ′-Fe4N thin films deposited on (a) MgO(100) substrate,

(b) SrTiO3(100) substrate, and (c) LaAlO3(100) substrate

-10000 -5000 0 5000 10000-2

0

2

me

mu

Applied Field (Oe)

(a)-Fe4N on MgO(100)

-10000 -5000 0 5000 10000-2

0

2(b)-Fe4N on SrTiO3(100)

me

mu

Applied Field (Oe)

-10000 -5000 0 5000 10000-2

0

2(c)-Fe4N on LaAlO3(100)

me

mu

Applied Field (Oe)


98

Figure 4.30 (a) Ms plotted against FWHM for the samples annealed for 10 – 40 minutes,

(b) Ms plotted against % lattice mismatch for the samples deposited on three

different substrates

0.3 0.6 0.9 1.2 1.51050

1200

1350

1500

1650

1800

1950

Ms (

em

u/c

c)

FWHM ()

0 3 6 9 121050

1200

1350

1500

1650

1800

1950LaAlO3

SrTiO3

MgOM

s (

em

u/c

c)

% Lattice Mismatch

(b)(a)


99

Table 4.6 A table showing % lattice mismatches between substrates and the deposit

along with the Ms and μB/Fe values for the γ′-Fe4N films

Sr.

No.

Substrate Lattice

constant

(Å)

Lattice

constant of γ′-

Fe4N

% Lattice

mismatch

Ms

(emu/cc)

±20

μB/Fe

1. MgO(100) 4.2 3.79 11 % 1760 2.59

2. SrTiO3(100) 3.9 3.79 3 % 1880 2.76

3. LaAlO3(100) 3.79 3.79 0 % 1980 2.91


100

4.7 Magnetic Anisotropy

Prior to the utilization of vibrating sample magnetometer for the calculation of

saturation magnetization, the ferromagnetic behavior of the films was also examined

using magneto-optic microscope magnetometer (MOMM) [114]. Strong ferromagnetic

behavior of γ΄-Fe4N thin films was revealed by the quite squarish magnetic hysteresis

loops with an in-plane applied magnetic field as shown in the Figures 4.31 (a-d). The

loops showed the easy axis in <100> direction with distinct fourfold cubic anisotropy,

characteristic of well-ordered single crystal epitaxial texture.

Angle dependent magnetic hysteresis loops were traced with 15 degree steps, in

order to find the direction of magnetic anisotropy for the sample deposited at 450 ˚C and

in-situ annealed for 30 minutes as shown in the Figure 4.32. The results revealed that

<100> is the easy axis of magnetization and <110> is the hard axis in the plane as the

loops showed the easy axis in <100> direction with distinct fourfold cubic anisotropy,

which is also characteristic of well-ordered single crystal epitaxial texture [120].


101

Figure 4.31 Hysteresis loops obtained using MOMM for the sample annealed for

(a) 10 min. (b) 20 min. (c) 30 min. and (d) 40 min.

-2 -1 0 1 2

80

82

84

Ke

rr I

nte

nsity (

a.u

.)

Applied Field (a.u.)

(a)

-2 -1 0 1 282

84

86

Ke

rr I

nte

nsity (

a.u

.)


(b)

-2 -1 0 1 2

80

82

84

Ke

rr I

nte

nsity (

a.u

.)


(c)

-2 -1 0 1 2

80

82

84

Ke

rr I

nte

nsity (

a.u

.)


(d)


102

Figure 4.32 Angle dependent hysteresis loops obtained using MOMM for the sample


-1.0 -0.5 0.0 0.5 1.094

96

98

Kerr

Inte

nsity (

a.u

.)


0 degree

-1.0 -0.5 0.0 0.5 1.093

94

95

96

15 degree

Kerr

Inte

nsity (

a.u

.)


-1.0 -0.5 0.0 0.5 1.092

93

94

95

30 degree

Kerr

Inte

nsity (

a.u

.)


-1.0 -0.5 0.0 0.5 1.092

93

94

95

96

45 degree

Kerr

Inte

nsity (

a.u

.)


-1.0 -0.5 0.0 0.5 1.093

94

95

96

60 degree

Kerr

Inte

nsity (

a.u

.)


-1.0 -0.5 0.0 0.5 1.092

93

94

95

96

75 degree

Kerr

Inte

nsity (

a.u

.)


-1.0 -0.5 0.0 0.5 1.095

96

97

98

90 degree

Kerr

Inte

nsity (

a.u

.)



103

4.8 Static Magnetic Domains

Strong ferromagnetic behavior was also verified with the help of magnetic force

microscope (MFM) for the sample deposited at 450 ºC and in-situ annealed for

30 minutes. MFM is a very powerful technique for studying static magnetic domain

structure of ferromagnetic materials as it allows imaging magnetic domains with a high

spatial resolution [121]. In MFM, the magnetic contrast is obtained through detecting the

force gradient between a ferromagnetic tip and the magnetic sample by amplitude, phase,

or frequency detection techniques. One distinguishing characteristic of imaging by MFM

is that it is compatible with external magnetic fields which make it suitable for

identifying magnetic domain structures [122].

Utilizing this technique, bow-tie-shaped static magnetic domain structures were

observed as shown in Figure 4.33. The dark-colored area of the domains indicates that tip

is repulsive to the domain area and light-colored area indicates it is attractive [123].

Traditionally this type of domain structure is expected from the ferromagnetic materials

with perpendicular magnetic anisotropy (PMA).

The (100) surface of γ′-Fe4N film contains two easy directions, and the behavior

of the magnetically pristine sample is characterized by a preference for 90° domain walls.

The MFM image shows that the magnetic domains in the film are in the micrometer

range. The directions of magnetization of almost any two adjacent domains are aligned

along the easy axes and are perpendicular to each other. The divergence of local

magnetization at the domain boundary produces negative and positive magnetic charge

that appears as bright and dark contrast along the domain walls. Magnetic domains

imaged during magnetization reversal with the field applied along the easy axes of the

film have demonstrated that switching processes occur mostly via domain wall motion

[43].


104

Figure 4.33 Magnetic domain images captured using MFM for the γ′-Fe4N film deposited

at 450 ˚C and in-situ annealed for 30 minutes


105

4.9 TEM-Microstructure

The Transmission Electron Microscopy (TEM) combined with Lorentz

Microscopy and Electron Holography is considered quite a useful technique tool for

characterizing high density magnetic thin films. Figure 4.34 shows the micrograph of γ΄-

Fe4N films prepared at 450 ºC obtained with help of transmission electron microscope.

As compared with the micrograph obtained for pure iron film deposited on MgO(100)

under the same conditions (as shown in figure 4.35), a very dense and uniform

microstructure, having grains in the range of 60-70 nm, was observed for γ΄-Fe4N films

in-situ annealed for 30 minutes. Such a dense and uniform structure suggests the strong

exchange interaction among the magnetic domains resulting in the enhancement of

magnetic moment in γ΄-Fe4N thin films.

4.10 Lorentz Microscopy and Electron Holography

An electron wave after transmission through the specimen interferes with a

coherent reference wave and the corresponding interference pattern is recorded and the

phase shift can be recovered by mathematical equations. Lorentz microscopy is basically

meant for converting the phase information of the exit wave into an intensity information

of the detected wave front. Eventually, Electron holography reconstructs the wave field

which exits from a specimen both in amplitude and phase [124].

Figure 4.36 (a-b) shows the Lorentz image and the Figure 4.37 (a-b) shows the

electron holograms at different magnifications obtained for γ΄-Fe4N sample prepared at

450 ºC and in-situ annealed for 30 minutes. In the Lorentz microscope image, a white line

is observed as indicated by an arrow. This boundary is considered to be a magnetic

domain wall [125]. The presence of cyclic holographic vertices depicting strong magnetic

flux is a direct evidence of strong exchange coupling among the magnetic domains which

in turn resulted in the enhancement of magnetic moment for γ΄-Fe4N film. The broad

fringes arise due to the change in specimen thickness around the edge.


106

Figure 4.34 Micro-structural images obtained using TEM for the sample deposited at

450 ˚C and in-situ annealed for 30 minutes

Figure 4.35 Micro-structural images obtained using TEM for the pure iron deposited on

MgO(100)


107

(a)

(a)

(b)

Figure 4.36 Micro-structural images obtained using TEM for the sample deposited at

450 ˚C and in-situ annealed for 30 minutes in (a) TEM mode, and

(b) Lorentz mode


108

(a)

(b)

Figure 4.37 Holographic images obtained using TEM for the sample deposited at 450 ˚C

and in- situ annealed for 30 minutes in (a) high magnification and (b) low

magnification


109

4.11 Low Temperature Magnetic Behavior

During the last few weeks of the present research, we had the opportunity to use a

superconducting quantum interference device (SQUID) for a few days. Availing this

opportunity, we tried to explore the low temperature magnetic behavior of γ′-Fe4N thin

films, using SQUID. Figures 4.38 and 4.39 show magnetic hysteresis loops obtained at

300 K, 250 K, 200 K, 150 K, 100 K, 50 K, and 10 K for the γ′-Fe4N film deposited on

MgO(100) substrate at 450 ˚C and in-situ annealed for 30 minutes. It is obvious that the

value of magnetic moment increased as the temperature decreased from room

temperature (RT) to 10 K. This behavior of ferromagnetic iron nitride thin films is in

agreement with Weiss theory of ferromagnetism [126]. According to this theory, the

spontaneous magnetization within a ferromagnetic domain increases as the temperature

decreases below RT. As the temperature approaches absolute zero, a perfect alignment of

magnetic moments occurs within the domain which consequently enhances the overall

magnetic moment of the film. Figure 4.40 shows the trend of magnetic moment for the γ′-

Fe4N film as the temperature decreases from RT to 10 K. As can be seen in the plot, an

almost linear behavior is obvious revealing increase of magnetic moment as the

temperature decreases. Near 10 K, a sharp rise in normalized moment reveals that

magnetic moment within domains tends to align in a perfect manner as the temperature

approaches absolute zero. The coercivity of the γ′-Fe4N film was observed to be 50 Oe at

RT and increased almost linearly with the decrease in temperature up to a value of

140 Oe at 10 K as shown in Figure 4.41.

4.12 Electrical Behavior of γ′-Fe4N Films

An AC resistance bridge externally attached with SQUID was employed to

investigate the electrical behavior of the γ′-Fe4N films as the temperature decreases from

RT. This arrangement utilizes the Four Point Probe method to measure the resistivity of

the films. Analytical grade Ag paste was used to make ohmic contacts. The sample was

placed in a dessicator for 48 hours in order to dry out the paste. Figure 4.42 shows the

resistivity response as the temperature decreases from RT. The RT value of resistivity

came out to be 55 μΩ cm which is less than reported by Nikolaev et al. [43]. A metallic


110

Figure 4.38 M-H loops obtained at 300 K, 250 K, 200 K, and 150 K for the sample

deposited at 450 ˚C and in-situ annealed for 30 minutes

-6000 -4000 -2000 0 2000 4000 6000

-0.0004

-0.0002

0.0000

0.0002

0.0004

em

u

Applied Field (Oe)

300K

-6000 -4000 -2000 0 2000 4000 6000

-0.0004

-0.0002

0.0000

0.0002

0.0004

em

u

Applied Field (Oe)

250K

-6000 -4000 -2000 0 2000 4000 6000

-0.0004

-0.0002

0.0000

0.0002

0.0004

em

u

Applied Field (Oe)

200 K

-6000 -4000 -2000 0 2000 4000 6000

-0.0004

-0.0002

0.0000

0.0002

0.0004

em

u

Applied Field (Oe)

150K


111

Figure 4.39 M-H loops obtained at 100 K, 50 K, and 10 K for the sample deposited at

450 ˚C and in-situ annealed for 30 minutes

-6000 -4000 -2000 0 2000 4000 6000

-0.0004

-0.0002

0.0000

0.0002

0.0004em

u

Applied Field (Oe)

100K

-6000 -4000 -2000 0 2000 4000 6000

-0.0004

-0.0002

0.0000

0.0002

0.0004

0.0006

em

u

A

50K

-6000 -4000 -2000 0 2000 4000 6000-0.0006

-0.0004

-0.0002

0.0000

0.0002

0.0004

0.0006

em

u

Applied Field (Oe)

10K


112

Figure 4.40 Normalized magnetization plotted against temperature for the

sample deposited at 450 ˚C and in-situ annealed for 30 minutes

0 50 100 150 200 250 3000.80

0.84

0.88

0.92

0.96

1.00

M/M

s

Temperature (K)


113

Figure 4.41 Coercivity plotted against Temperature for the γ′-Fe4N thin films

0 50 100 150 200 250 30040

60

80

100

120

140

Coerc

ivity (

Oe)

Temperature (K)


114

behavior of the γ′-Fe4N was witnessed as the resistivity decreased quite linearly with the

decrease of temperature. This value of RT resistivity of γ′-Fe4N films is quite optimum to

reduce eddy current losses and to extend the flat frequency response up to the high

frequency range, and make these films a strong candidate to be used as write head

materials [127]. The exact resistivity of the film could not be measured below 80 K due

to contact breakage which might be due to non-linear contraction of the film and the

contacts at low temperature. The extraplotation of the resistivity vs temperature curve

gives a projected value of resistivity of about 20 μΩ cm at 10 K.


115

Figure 4.42 Resistivity plotted against temperature for the the sample deposited at 450 ˚C

and in-situ annealed for 30 minutes

0 50 100 150 200 250 300

2.0x10-5

2.5x10-5

3.0x10-5

3.5x10-5

4.0x10-5

4.5x10-5

5.0x10-5

5.5x10-5

Resis

tivity (

cm

)

Temperature (K)

Chapter 5 Preparation and Characterization of α̋-Fe16N2 Thin Films

116

Preparation and Characterization of α″-Fe16N2

Thin Films

(Results and Discussions II)

Light elements like hydrogen, boron, carbon and nitrogen at the beginning of the

periodic table are small enough in their covalent or metallic state to interstitially place

themselves in the magnetic 3d lattice of iron. Oxygen, fluorine, and chlorine being

strongly electronegative, tend to form negative ions. Hydrogen is interstitially insoluble

in iron, while carbon can dissolve in both α- and γ-Fe to yield the most useful alloys

known to mankind. For the formation of nitrides, the usual approach is a gas-phase

interstitial modification at quite low temperatures (> 670 K) [28].

The nitrides of iron have been playing an important role in steel technology.

Particularly, nitriding is used to harden the iron surface and to passivate them against

oxidation. In addition, the magnetic properties of iron nitrides with small N contents, in

combination with their favorable corrosion and wear properties, have raised considerable

interest concerning their applications in magnetic data storage [128].

Among ferromagnetic phases of iron nitrides, α″-phase is one of the most

fantastic and mysterious materials in the field of magnetism owing to the prolonged

debate about its magnetic moment. In-spite of the confirmation of giant magnetic

moment in α′′-Fe16N2 thin films [129], it is still a controversial material, attracting much


117

attention of the researchers worldwide. The criticality started in 1972 [12], when a giant

value of saturation magnetization was reported in epitaxial α′′-Fe16N2 thin films prepared

by molecular beam epitaxy (MBE), after about two decades when this material was first

structurally discovered [21]. This report made a strong impact on the research fields of

magnetic materials and many research groups tried to clarify its magnetism. For the

purpose, different thin film processing techniques have been employed to synthesize this

compound but mostly led to the variable results. For instance, Sugita et al. [58] have

claimed reconfirming the giant magnetization of this phase but the work of Shoji et al.

[130] seems to dispute this claim. The contradiction might be due to the fact that α′′-

Fe16N2 exists only as a meta-stable phase and thermally less stable [21] as compared with

the other strong ferromagnetic phase of iron nitride. Efforts are continue worldwide to

probe into this dilemma. Therefore, the study of magnetic properties of α′′-Fe16N2 phase

is very important for both fundamental physical research as well as technological

applications.

5.1 Preparation of α′′-Fe16N2 Thin Films

The preparation of single phase epitaxial thin films of α″-Fe16N2 has always been

a challenging task for the researchers worldwide. The reason might be due to the fact that

it is metastable phase and thermally less stable [21].

In order to achieve the optimum conditions for single phase growth of α″-phase of

iron nitride, we deposited several series of thin films on MgO(100), and Si(100)

substrates employing different sets of parameters with varying nitrogen partial pressure,

substrate temperature, annealing time, dc sputtering power, and the film thickness. The

base pressure of the chamber was better than 2 x 10-6

Torr. Analytically pure Ar gas

(99.95% purity) was flown in to the chamber as a working gas. The substrates were

cleaned by chemical and ultra sonic treatment prior to deposition. The complete process

of substrate cleaning has been described in section 3.2.1. Utilizing optimized deposition

conditions, following series of samples (named as „5‟) was deposited, which eventually

led to the single phase epitaxial growth of α″-Fe16N2 films on single crystal Si(100)

substrates.


118

(5) Changing nitrogen partial pressure from 0.2 to 1.0 mTorr, at substrate temperature

of 200 ˚C, and in-situ annealing of 1 hour. The films were deposited at a

deposition rate of 0.88 Å/sec utilizing dc sputtering power of 60 W. The

thickness of each film was 40 nm. The complete set of deposition conditions is

shown in Table 5.1


119

Table 5.1: Nitrogen partial pressure dependent preparation of α″-Fe16N2, Series 5


PAr = 5.0 mT

Thickness = 40 nm



(˚C)

Annealing Time

(Minutes)

Substrate

Material

5(a) 0.2 200 1.0 Si(100)

5(b) 0.4 200 1.0 Si(100)

5(c) 0.6 200 1.0 Si(100)

5(d) 0.8 200 1.0 Si(100)

5(e) 1.0 200 1.0 Si(100)


120

5.2 Crystal Structure Determination

Figure 5.1 shows the XRD patterns of iron nitride thin films deposited on single

crystal Si(100) substrates with nitrogen partial pressure (2NP )varying from 0.2 to

1.0 mTorr. All the films were deposited at the substrate temperature of 200 ºC and in-situ

annealed for 1 hour. The sample deposited at 2NP = 0.2 mTorr showed relatively small x-

ray intensity peaks. The pattern revealed bi-phase iron nitride structure as the peaks

corresponding to two different phases, Fe3N, and α′′-Fe16N2 were observed as indicated in

the Figure 5.1 (b). All the remaining peaks were from Si substrate. For comparison, the

XRD pattern of Si(100) substrate used in this study is shown in Figure 5.1 (a).

The increase of 2NP to 0.4 mTorr resulted in the enhancement of peak intensity

and two relatively clear peaks of α′′-Fe16N2(220) at 2θ = 44.72º and that of γ′-Fe4N(200)

at 2θ = 47.91º were observed as shown in Figure 5.1 (c). Nearly the same bi-phase texture

was observed when the 2NP was increased to 0.6 m Torr, with the only difference that an

additional (110) peak of α′′-Fe16N2 was also seen at 2θ = 21.95º (Figure 5.1d). Single

phase epitaxial texture of α′′-Fe16N2 phase was observed for nitrogen partial pressure of

0.8 m Torr as the only (220) peak of the phase was observed at 2θ = 44.72º, as shown in

Figure 5.1 (e). This illustrated that 0.8 mTorr is optimum partial pressure of nitrogen for

the single phase epitaxial growth of α′′-Fe16N2 thin films as the increase in 2NP to a value

of 1.0 mTorr resulted again in the appearance of multi-phase iron nitride structure, as can

be seen in Figure 5.1 (f).


121

Figure 5.1 XRD patterns of (a) pure Si(100) substrate, iron nitride samples deposited on

Si(100) using nitrogen partial pressure of (b) 0.2 mTorr, (c) 0.4 mTorr,

(d) 0.6 mTorr, (e) 0.8 mTorr, and (f) 1.0 mTorr

20 30 40 50 60 70

Fe16N2

Fe4N

Fe3N

Si Substrate

2()

(a)

(b)

(c)

(d)

(e)

(f)


122

5.3 Structural Morphology of Surface

The surface morphology of the iron nitride films was investigated using scanning

electron microscopy as it helps us to view the surface in an optical-like form. The image

is formed by the secondary electrons emitted from the surface. The intensity and angle of

emission of the electrons both depend on the surface topography of the iron nitride

samples [131]. Figures 5.2 – 5.5 show the surface images of samples prepared utilizing

nitrogen partial pressure of 0.4 – 1.0 mTorr, respectively. All the samples were prepared

at the same substrate temperature and annealed for the same period of time, therefore the

surface images of all the samples showed almost similar structure with the grain size

ranging from 15 to 20 nm. The only difference seen was in the surface image of the

sample prepared at nitrogen partial pressure of 0.8 mTorr, where some grain clusters

were observed consisting of three to four individual grains. Nevertheless, it is evident that

a mere change in nitrogen partial pressure, with other parameters undisturbed, does not

affect the grain size of the films.

The change in the lateral grain size of the thin film samples is mostly governed by

the change in super-saturation and the surface mobility of the adsorbed species. The

decrease in super-saturation or increase in surface mobility results in the increase of grain

size. Both these factors are directly related with substrate and source temperature [132].

If these conditions remain undisturbed, the grain size remains nearly the same.

5.4 RMS Roughness of Surface

Surface roughness is among one of the important parameters to judge the

structural quality of a thin film surface. Figure 5.6 shows the graph plotted between

nitrogen partial pressure and the root mean square (RMS) value of surface roughness

obtained using atomic force microscope. The RMS value of surface roughness for the

series of samples deposited at nitrogen partial pressures of 0.2 - 1.0 mTorr was observed

to be almost same, lying in the range of 0.51 (± 0.02) nm, except for the sample deposited

at a nitrogen partial pressure of 0.8 mTorr which showed an RMS value of 0.38 (±0.02)

nm for the surface roughness. This result seems to be consistent as was observed in SEM

surface image for the sample that showed the coalescence of individual grains at certain


123

Figure 5.2 Microstructural image obtained by SEM for the iron nitride sample deposited

using 2NP = 0.4 m Torr




124






125

Figure 5.6 Effect of nitrogen partial pressure on the RMS value of surface roughness

0.2 0.4 0.6 0.8 1.0

0.3

0.4

0.5

0.6

RM

S R

ou

gh

ne

ss (

nm

)

Nitrogen Partial Pressure (mTorr)


126

places. However, these are relatively large values of surface roughness as compared with

the values discussed in section 4.4 for γ′-Fe4N films (deposited at 450 ˚C). The reason

might be the fact that the films under discussion are deposited at relatively low

temperature of 200 ˚C. Therefore, less surface mobility has resulted in relatively large

values of RMS roughness.

The surface roughness and the microstructure of the films are mostly affected by

the strains developed due to the thermal expansion caused by the lattice mismatch

between a film and its substrate [133]. Furthermore, during growth process, under the

conditions like low nucleation barrier and high super-saturation, the initial nucleation

density is high which results in fine-grained and smooth deposit of the films. On the

contrary, if the nucleation barrier is large and the super-saturation is low, course-grained

rough films are formed. Generally, high surface mobility increases the surface

smoothness of the films by filling in the concavities.

5.5 Magnetic Domain Structure

Figure 5.7 shows the magnetic domain structure and Figure 5.8 shows the

corresponding surface topography of the α″-Fe16N2 thin film obtained using magnetic

force microscopy. The film was prepared using nitrogen partial pressure of 0.8 mTorr at a

substrate temperature of 200 ˚C and in-situ annealed for 1 hour. The magnetic contrast in

magnetic force microscopy is obtained through detecting the force gradient between a

ferromagnetic tip and the magnetic sample by amplitude, phase, or frequency detection

techniques.

The dark and bright contrast depicts the presence of high moment magnetic

domains, although sharp boundaries could not be observed among the domains. The

reason might be the less smooth surface of the film as obvious in the surface topography

achieved for the sample compared to γ′-Fe4N films discussed in section 4.8.

5.6 Saturation Magnetization

The saturation magnetization of iron nitride thin films deposited on Si(100)

substrates under varying nitrogen partial pressure was also determined using vibrating

sample magnetometery. Film thickness being very critical in this study for the accurate


127

Figure 5.7 Magnetic domain structure obtained using MFM for iron nitride sample

deposited using nitrogen partial pressure of 0.8 mTorr

Figure 5.8 Surface topography for iron nitride sample deposited using nitrogen partial

pressure of 0.8 mTorr


128

volume calculation was determined by the cross-sectional images obtained using

scanning electron microscope.

The magnetic-hysteresis loops of the films obtained using VSM are shown in

Figure 5.9. The value of saturation magnetization (Ms) of iron nitride thin films plotted as

a function of 2NP is shown in the Figure 5.10. Minimum Ms value of 1350±20 emu/cc

was observed for the sample prepared using 2NP of 0.2 mTorr. This might be attributed

to the presence of Fe3N phase (Figure 5.1 b) of iron nitride which exhibits least

ferromagnetic behavior as compared with other ferromagnetic iron nitride phases [134].

When the 2NP was increased to 0.4 mTorr, the XRD pattern (Figure 5.1 c) showed the

presence of γ′-Fe4N along with α′′-Fe16N2, the two strong ferromagnetic phases of Fe-N

system, so the Ms increased to 1390±20 emu/cc.

Further increase in 2NP to 0.6 mTorr, resulted in the increase in α′′-Fe16N2

concentration, as was revealed by the XRD pattern of the sample shown in Figure 5.1 (d)

where in addition to (220) peak of α′′-Fe16N2, an additional (110) peak of α′′-Fe16N2 was

also seen. As a result the corresponding value of saturation magnetization increased to

1670±20 emu/cc. This increase in Ms was obviously due to the increased concentration of

α′′-Fe16N2 phase in the film. When the 2NP was increased gradually up to a value of 0.8

mTorr, the Ms reached to a maximum value of 1800±20 emu/cc. XRD pattern (Figure 5.1

e) of this sample illustrated single phase epitaxial growth of α′′-Fe16N2 films, so a high Ms

could justifiably be expected from this sample [135,50]. The influence of 2NP on the

phase formation and the corresponding Ms is presented in Table 5.2.

This result is consistent as well with the observed value of RMS roughness for

this sample. If nitrogen partial pressure is plotted against Ms and RMS value of surface

roughness as shown in the Figure 5.11, it is observed that the film prepared using2NP

value of 0.8 mTorr showed minimum RMS value of surface roughness and maximum

value of saturation magnetization which is consistent as well with our previous findings

for γ′-Fe4N films [50]. However this maximum Ms value of 1800±20 emu/cc observed for

epitaxial α′′-Fe16N2 film is about 14% less than the giant value of saturation

magnetization reported by Kim et al. for this material [12].


129

Figure 5.9 M – H loops obtained using VSM for the iron nitride samples deposited

using nitrogen partial pressure of (a) 0.2 m Torr, (b) 0.4 m Torr,

(c) 0.6 m Torr, (d) 0.8 m Torr, and (e) 1.0 m Torr

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(a)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(b)

-15000 -7500 0 7500 15000

-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(c)

-15000 -7500 0 7500 15000

-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(d)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(e)


130

Figure 5.10 Effect of nitrogen partial pressure on saturation magnetization

0.2 0.4 0.6 0.8 1.01300

1400

1500

1600

1700

1800

Ms (

em

u/c

c)

Partial Pressure of N2 (m Torr)


131

Table 5.2 The Influence of nitrogen partial pressure on phase formation and the

corresponding Ms value

Sr. No.

2NP ( mTorr) Phases developed Ms (emu/cc) ±20

1.

0.2 α″-Fe16N2, γ′-Fe4N, Fe3N 1350

2.

0.4 α″-Fe16N2, γ′-Fe4N 1390

3.

0.6 α″-Fe16N2, γ′-Fe4N 1670

4.

0.8 α″-Fe16N2 1800

5.

1.0 α″-Fe16N2, γ′-Fe4N, Fe3N 1475


132

Figure 5.11 Effect of nitrogen partial pressure on saturation magnetization

and RMS value of surface roughness for iron nitride thin films

0.2 0.4 0.6 0.8 1.0

1300

1400

1500

1600

1700

1800

1900

Nitrogen Partial Pressure (mTorr)

Ms (

em

u/c

c)

0.3

0.4

0.5

0.6

RM

S R

oug

hn

ess (

nm

)


133

The sample prepared at 2NP = 1.0 mTorr showed a decrease in Ms which is again due to

the mixed phase iron nitride film and the presence of less ferromagnetic phase (Fe3N) of

iron nitride as revealed by the XRD pattern of the sample shown in Figure 5.1 (f).

5.7 Motivation for Studying the α′′-(Fe,X)16N2 System

The contradictory results about the giant value of saturation magnetization might

be due to the fact that α′′-Fe16N2 exists only as a meta-stable and thermally less stable

phase [21] as compared to the second strong ferromagnetic phase of iron nitride [44,50].

Jack [71] has suggested to improve the thermal stability employing both interstitial and

substitutional alloying elements to develop α′′-(Fe,X)16N2 (X = metal) phases, with the

aim of incorporating them into thin film heads for high-density disk drives. Therefore, the

study of magnetic properties of α′′-Fe16N2 phase and the effect of alloying addition on its

magnetic properties is very important as far as the technological applications in recording

media is concerned.

Presently, permalloys are mostly being used for the head materials due their soft

magnetic properties, ease of fabrication and low magnetostriction. However, there is

growing realization that high density recording cannot be accomplished without

permalloys being replaced by other magnetic materials such as FeXN [136]. This is the

driving force behind the research for materials with high saturation magnetic moment

such as iron nitrides and the subsequent addition of alloying elements.

In this study, we have investigated the nitrogen partial pressure dependent

preparation of α″-phase of iron nitride and the subsequent effect of Co, Pt, and Cr

additions on the saturation magnetization of single phase epitaxial α′′-Fe16N2 films. The

reasons for Co additions as a third element, are twofold: first, Fe maintains a body-centre

cubic (bcc) structure even after alloying with the third element which is essential as α′′-

phase has as analogous bct structure, and secondly, the third element does not undergo

preferential nitridation as its chemical affinity for nitrogen is weaker or comparable to

that of Fe [137]. Co meets these two requirements enabling itself a proper choice for

alloying addition. Pt addition as an alloying element could induce relatively large in-

plane coercivity required for potential applications in recording media [138]. Addition of

Cr in iron nitride might be suitable as well, as Cr has less than 1% lattice mismatch with

Fe and both the elements have bcc structure [139].


134

5.8 FeCoN Alloy System

5.8.1 Preparation

Utilizing the deposition conditions of 2NP = 0.8 mTorr, 200 ºC substrate

temperature and in-situ annealing of 1 hour, as were optimized for single phase epitaxial

growth of α′′-Fe16N2 thin films on Si(100), the FeCoN films were prepared. For this

purpose, Co chips (size: 4 × 4 mm) were placed on the iron target having diameter of

50 mm. The Co atomic concentration percent (at%) in FeN films was controlled by

varying the number of Co chips from 1 to 4 during each deposition. All other deposition

conditions like base pressure, working gas pressure, and sputtering power were same as

those used for epitaxial film of α′′-Fe16N2.

5.8.2 Co Composition Determination

The Co atomic concentration percent in iron nitride was evaluated using Auger

Electron Spectroscopy (AES). The technique is based on the measurement of the energy

and intensity of the Auger electron signal emitted from the atoms located within a few

nanometers of the surface. These electrons are excited by a beam of incident electrons.

The energies of the emitted auger electrons are characteristics of the atoms emitting them,

while the magnitude of the AES signal is related directly to the abundance of the ad-

atoms under consideration and thus serve to fingerprint the elemental composition [140].

In this series of samples, the no. of chips was changed from 1 to 4 in order to get

increased concentration of Co in iron nitride matrix. The quantitative estimation of Co

at% was evaluated from Survey Scan Data obtained using AES, as shown in Figure 5.12.

Atomic concentration Cx of an element „x‟ may be estimated by

SI

SIC xx

x

………………………..…….(5.1)

where, Ix is the peak to peak height of the Auger electron signal for the element „x‟

divided by the proper scale sensitivity and Sx is the sensitivity factor. The AES spectra

were obtained at a base pressure of 8.0 × 10-10

Torr and a working gas (Ar: for pre-

sputtering) pressure of 9.0 × 10-8

Torr. Each sample was pre-sputtered for 24 seconds to


135

remove any contamination and oxide layers up to few nanometers from the sample

surface. The Co at% was increased almost linearly by increasing the number of Co chips.

The elemental composition of Co in iron nitride is shown in the Table 5.3.

5.8.3 Effect of Co Addition on Ms

Figure 5.13(a-e) shows the M-H loops of the FeCoN alloy films having Co at% of

0, 4.37%, 9.68%, 12.86%, and 16%, respectively and Figure 5.14 shows the plot of Co

at% in iron nitride vs Ms and the number of Co chips used. It was observed that minimum

value of saturation magnetization (1400±20 emu/cc) was achieved when the Co at% was

4.37% corresponding to 1 Co chip. Afterwards, the Ms increased and reached a maximum

value of 1660±20 emu/cc as the Co concentration reached to a value of 12.86 at% by

changing the number of Co chips from 1 to 3. This value of Ms is a little higher than

observed by Takahashi et al. [141] for α′-(Fe-Co)16N2 phase. At 16 at% of Co,

corresponding to 4 Co chips, the Ms decreased again to a value of 1515±20 emu/cc. From

this we deduced that the alloying addition of Co in α′′-Fe16N2 phase could be favorable as

far as stability is concerned, with some compromise on saturation magnetization as

compared with pure α′′-Fe16N2 phase. Sample preparation under ultra high vacuum

conditions followed by compositional and magnetic analysis employing more accurate

and precise instruments might be helpful in achieving even better results.


136

Figure 5.12 AES scans obtained for FeCoN films deposited by varying Co chips

200 400 600 800 1000

Kinetic Energy (eV)

No chip

2 chips

1 chip

3 chips

4 chips

Iron

Cobalt

Nitrogen

Impurity


137

Table 5.3 The elemental composition of Co in iron nitride corresponding to

number of Co chips

Sample

No

No. of Co

chips

% of Fe % of N % of Co

B1

0 88.97 11.02 0

B2

1 84.99 10.63 4.37

B3

2 79.89 10.42 9.68

B4

3 77.99 9.14 12.86

B5

4 76.20 7.78 16.00


138

Figure 5.13 M – H loops obtained by VSM for iron nitride films deposited using (a) No

Co chip, (b) 1 Co chip, (c) 2 Co chips, (d) 3 Co chips, and (e) 4 Co chips

-15000 -7500 0 7500 15000

-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(a)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(b)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(c)

-10000 -5000 0 5000 10000

-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(d)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(f)


139

Figure 5.14 Effect of no. of Co chips on at% Co in α″-Fe16N2 and

on saturation magnetization

0 4 8 12 16

1400

1500

1600

1700

1800

Co at% in -Fe16N2

Ms (

em

u/c

c)

0

2

4

No

. o

f C

o c

hip

s


140

5.9 FePtN Alloy System

5.9.1 Preparation



growth of α′′-Fe16N2 thin films on Si(100), the FePtN films were prepared. Due to the fact

that non-ferromagnetic metals are relatively easy to sputter using magnetron sputtering,

so the Pt a non-ferromagnetic metal exhibits a high deposition rate. Therefore, the size of

the Pt chips used was 2 × 4 mm, half of that as was used for Co chips. The Pt chips of

analytical purity (99.95% pure) were placed on the iron target having diameter of 50 mm.

The Pt atomic concentration percent in FeN films was controlled by varying the number

of Pt chips from 1 to 4 during each deposition. All other deposition conditions like base

pressure, working gas pressure, and sputtering power were same as used for epitaxial film

of α′′-Fe16N2.

5.9.2 Pt Composition Determination

The Pt atomic concentration percent in iron nitride was evaluated using Auger

Electron Microscope (AES) as shown in Figure 5.15. Although, the size of Pt chips used

was relatively small, even then high concentration of Pt was observed in the FePtN films

as evaluated by the Survey Scan Data obtained by AES spectra using Eq. 5.1. All the

conditions for compositional analysis were the same as described in section 5.8.2. The Pt

at% was increased almost linearly by increasing the number of Pt chips from 1 to 4. The

elemental composition of Pt in iron nitride is shown in the Table 5.4.

5.9.3 Effect of Pt Addition on Ms

Figure 5.16(a-e) shows the M-H loops of the FePtN alloy films having Pt at% of 0,

16.48%, 22.54%, 27.46%, and 35.47%, respectively and Figure 5.17 shows the plot of Pt

at% in iron nitride vs Ms and the number of Pt chips used. It was observed that minimum

value of saturation magnetization (1300±20 emu/cc) was achieved when the Pt at% was

16.48% corresponding to 1 Pt chip. Afterwards, the Ms increased and reached a

maximum value of 1510±20 emu/cc as the Pt concentration reached to a value of 27.46

at% by changing the number of Pt chips from 1 to 3. Additionally, this sample exhibited


141

maximum value of coercivity (≈ 1.2 kOe) in the FePtN series of samples. In literature, no

report has been found stating the effect of Pt addition on saturation magnetization of iron

nitride films. At 35.47 at% of Pt, corresponding to 4 Pt chips, the Ms remained almost

unchanged as was seen in the sample prepared using three Pt chips. More systematic

investigations are needed to explore FePtN alloy phases to make this system compatible

for high density recording media applications that require high Ms with relatively high

coercivity [142].


142

Figure 5.15 AES scans obtained for FePtN films deposited by varying Pt chips

500 1000 1500 2000

Iron

Platinum

Nitrogen

Impurity

Kinetic Energy (eV)

1 Pt chip

2 Pt chips

3 Pt chips

No Pt chip

4 Pt chips


143

Table 5.4 The elemental composition of Pt in iron nitride corresponding to number

of Pt chips

Sample

No

No. of Pt

chips

% of Fe % of N % of Pt

C1

0 88.97 11.02 0

C2

1 73.38 10.13 16.48

C3

2 67.96 9.49 22.54

C4

3 64.07 8.460 27.46

C5

4 57.39 7.13 35.47


144


Pt chip, (b) 1 Pt chip, (c) 2 Pt chips, (d) 3 Pt chips, and (e) 4 Pt chips

-15000 -7500 0 7500 15000

-4

-2

0

2

4m

em

u (

a.u

.)

Applied Field (H)

(a)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(b)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Fiels (H)

(c)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u)

Applied Field (H)

(d)

-15000 -7500 0 7500 15000-4

-2

0

2

4

me

mu

(a

.u.)

Applied Field (H)

(e0


145

Figure 5.17 Effect of no. of Pt chips on at% Pt in α″-Fe16N2 and


0 10 20 30 40

1300

1400

1500

1600

1700

1800

Pt at% in -Fe16N2

Ms(e

mu/c

c)

0

2

4

No.

of

Pt

chip

s


146

5.10 FeCrN Alloy System

5.10.1 Preparation



growth of α′′-Fe16N2 thin films on Si(100), the FeCrN films were prepared. For this

purpose, Cr chips (size: 4 × 4 mm) were placed on the iron target having diameter of

50 mm. The Cr atomic concentration percent (at%) in FeN films was controlled by

varying the number of Cr chips from 1 to 4 during each deposition. All other deposition

conditions like base pressure, working gas pressure, and sputtering power were same as

used for epitaxial thin film of α′′-Fe16N2.

5.10.2 Cr Composition Determination

The Cr atomic concentration percent in iron nitride was also evaluated using

Auger Electron Microscope (AES) as shown in Figure 5.18. Eq. 5.1 was used in order to

calculate the elemental atomic percent using the Survey Scan Data obtained from AES

spectra. All the conditions for compositional analysis were the same as described in

section 5.8.2. The Cr at% was increased almost linearly by increasing the number of Cr

chips from 1 to 4. The elemental composition of Cr in iron nitride is shown in the Table

5.5.

5.10.3 Effect of Cr Addition on Ms

Figure 5.19 (a-e) shows the M-H loops of the FeCrN alloy films having Cr at% of

0, 6.99%, 12.00%, 16.06%, and 20.37%, respectively and Figure 5.20 shows the plot of

Cr at% in iron nitride vs Ms and the number of Cr chips used. It was observed that

maximum value of saturation magnetization observed in this series was 1350±20 emu/cc

at 6.99 at% of Cr corresponding to 1 Cr chip. The Ms decreased by the increase in Cr at%

concentration as the number of Cr chips was further increased. This trend might be

attributed to the fact that ferromagnetic-antiferromagnetic (FM-AFM) coupling between

Fe and Cr atoms increased as the Cr contents increased [143] in iron nitride. Therefore,

Cr could not favor its addition at all, in iron nitride to affect the Ms positively.


147

Figure 5.18 AES scans obtained for FeCrN films deposited by varying Cr chips

200 400 600 800 1000

Iron

Cromium

Nitrogen

Impurity

Kinetic Energy (eV)


148

Table 5.5 The elemental composition of Cr in iron nitride corresponding to number

of Cr chips

Sample

No

No. of Cr

chips

% of Fe % of N % of Cr

D1

0 88.97 11.02 0

D2

1 81.44 11.56 6.99

D3

2 79.60 11.38 12.00

D4

3 72.35 11.58 16.06

D5

4 66.41 11.76 20.37


149


Cr chip, (b) 1 Cr chip, (c) 2 Cr chips, (d) 3 Cr chips, and (e) 4 Cr chips

-15000 -7500 0 7500 15000

-4

-2

0

2

4m

em

u (

a.u

.)

Applied Field (H)

(a)

-15000 -7500 0 7500 15000

-2

0

2

me

mu

(a

.u.)

Applied Field (H)

(b)

-15000 -7500 0 7500 15000

-2

0

2

me

mu

(a

.u.)

Applied Field (H)

(c)

-15000 -7500 0 7500 15000-2

-1

0

1

2

me

mu

(a

.u.)

Applied Field (H)

(d)

-15000 -7500 0 7500 15000-1

0

1

me

mu

(a

.u.)

Applied Field (H)

(e)


150

Figure 5.20 Effect of no. of Cr chips on at% Cr in α″-Fe16N2 and


0 4 8 12 16 20300

600

900

1200

1500

1800

Cr at% in -Fe16N2

Ms (

em

u/c

c)

0

2

4

No

. o

f C

r ch

ips

Chapter 6 Power Law Scaling Behavior in 2D Ferromagnetic Films

151

Power Law Scaling Behavior in 2D Ferromagnetic Films

(Results and Discussions III)

Magnetic domain structure plays a vital role in assigning the soft magnetic

properties to a ferromagnetic material. In ferromagnetic thin films, magnetization reversal

mechanism is a subject of great interest. From both the fundamental and technological

aspects, the key interest lies in the Barkhausen avalanches during the magnetization

reversal process. Through a statistical analysis, it has been shown that the obvious

randomness of the jump size during magnetization reversal in iron nitride films, actually

follows a power law behavior which is also independent of the number of defects in the

films.

6.1 Magnetic Domains

In a ferromagnetic material, all the atomic moments in a single magnetic domain

are aligned parallel, so that one would expect the material to be magnetically saturated all

the times. The demagnetized or partially magnetized state is achieved by the material

being sub-divided into small regions, called „domains‟. Each one of them may be

magnetized locally to saturation but adding vectorially to a lower value, or to zero. In

other words, magnetic domains are the regions in a ferromagnetic material within which

the direction of magnetization is largely uniform. Once the domains are formed, the

orientation of magnetization in each domain and the domain size are determined by

magnetostatic, crystal anisotropy, magnetoelastic, and domain wall energy. The entire

domain structure calculations involve minimization of the appropriately selected energies

[144].


152

6.2 Domain Wall Motion

A domain wall is a region in which the local magnetization rotates from one easy

direction to another. In a cubic material, with k > 0 (positive anisotropy constant), the

easy directions are either 90˚ or 180˚ apart, and the domain walls are classified as 90˚ or

180˚ walls. Figure 6.1 (a) shows how magnetostatic energy depends on the presence of

domains. The domain formation in a saturated magnetic material is driven by the thrust to

lower the magnetostatic (MS) energy of the single domain state, as shown in left side of

Figure 6.1 (a). Introduction of 180˚ domain walls reduces the MS energy but raises the

wall energy. The 90˚ closure domains eliminate MS energy but increase the anisotropy

energy in uniaxial materials. The internal structure of a 180˚ domain is shown in Figure

6.1 (b). In most ferri- and ferromagnetic materials, the domain wall thickness is several

hundred atomic spacings. This means that there is some interaction between a domain

wall and a single impurity atom, a lattice vacancy, or an isolated dislocation [145].

Among magnetic domain walls, two types are the most important: Bloch and Neel

domain walls. Another type which is intermediate between Bloch and Neel walls, called

the cross-tie wall, has also been witnessed. The Bloch wall is preferable in bulk materials

where the spins rotate in the plane parallel to the wall plane (Figure 6.2-a). However, in

thin films, the Bloch wall induces surface charges by its stray field. When the film

thickness becomes smaller than the wall width, the Neel wall becomes more favorable. In

a Neel wall, spins rotate in the film plane as shown in the Figure 6.2 (b).

6.3 Magnetization Reversal

The reversal of magnetization or the switching behavior of magnetic domains in

ferromagnetic thin films is the subject of great interest from both fundamental and

technological aspects. This mechanism represents the process that leads to a 180˚

reorientation of the magnetization vector with respect to its initial direction, moving from

one stable orientation to the opposite one along the easy axis [146]. As for technological

applications are concerned, this is one of the most important process in magnetization

which is linked to the magnetic data storage process currently being used, for instance, in

the hard disc drive [147].


153

Figure 6.1(a), (a) Domain formation in a saturated magnetic material driven by the

magnetostatic (MS) energy in the single domain state, Introduction of 180o domain walls

reduces the MS energy but raises the wall energy creating (b) double domain state, (c)

multi-domain state, and (d) 90o closure domains eliminate MS energy but increase

anisotropy energy in a uniaxial material

Figure 6.1 (b) The internal structure of 180˚ domains


154

Figure 6.2 The rotation of the magnetization vector in the (a) Bloch wall, and

(b) Neel wall


155

Following are the possible ways to reverse the magnetization of a metallic

ferromagnetic material.

6.3.1 Magnetization Reversal by Spin Injection

In this process, a charge current of random spin polarization is injected through a

pillar composed of different thin layers. A ferromagnetic layer with fixed magnetization

direction, the polarizer, creates a spin polarized current as the injected current is passing

through this layer. Next, the spin polarized current enters the free layer and generates a

torque on the magnetization of this second ferromagnetic film. At sufficiently high

current densities the spin current is capable of switching the magnetization of free layer.

By reversing the direction of spin injected current, the magnetization can be switched

back to its initial direction.

6.3.2 All-optical Switching

All-optical switching is a brand new technique employed experimentally quite recently

[148] for the reversal of magnetization. The technique refers to a method for

magnetization reversal in a ferromagnetic material simply by circularly polarized light

where the magnetization direction is controlled by the light helicity. In particular, the

direction of the angular momentum of the photons would set the magnetization direction.

In-fact, this process could be seen as similar to magnetization reversal by spin injection

with the only difference that the angular momentum is supplied by the circularly

polarized photons instead of the polarized electrons.

6.3.3 Magnetization Reversal in an Applied Field

The conventional way to switch the magnetization is by an external magnetic

field. In order to switch the magnetization, in general, the magnetic field is applied

opposite to the initial magnetization orientation. The process of reversal takes place on a

nanosecond time scale and is dominated in thin films either by domain nucleation process

or by domain wall propagation.

Depending upon the applied field strength relative to the coercive field, there are

different reversal mechanisms. If an applied field is lower than the coercivity field,

magnetization reversal takes place by thermal activation of both nucleation and domain


156

wall motion process exhibiting exponential behavior in this regime [149]. When an

applied field becomes larger than the coercivity field then the magnetization is mainly

carried out by viscous domain wall motion. In this region, the domain wall velocity varies

linearly with an applied field [150]. For the much larger applied fields than the coercivity

field, the magnetization reversal is carried out via spin precessional motion governed by

Landau-Lifshitz-Gilbert equation [151].

6.4 The Barkhausen Effect

The interest in magnetization reversal has grown quite recently for an applied

field near the coercivity field. This is called the critical regime as the random avalanche

behavior in magnetization reversal process takes place across this region. In this regime,

the domain walls do not move smoothly and continuously, but in a series of local

displacements having discrete and jerky jumps as shown in Figure 6.3. This effect is

known as the “Barkhausen effect”, first discovered in 1919 [152] and results from local

pinning of walls by magnetostatic and magnetoelastic forces. These Barkhausen jumps

are the main cause of the noise heard in magnetic devices involving magnetization

reversal. We can practically "hear" the switching of magnetic domains by amplifying the

currents induced in a coil that surrounds the ferromagnetic material. The Barkhausen

effect has provided a direct evidence for the existence of ferromagnetic domains, which

were previously postulated theoretically.

6.5 The Critical Scaling Behavior

The main reason why interest has been revived in this classical subject is mainly

motivated by a fundamental question of whether there is any simple law governing this

seemingly random avalanche of magnetic domains during reversal phenomenon. It seems

to be an essential issue concerning our curiosity for understanding fundamental

magnetism, as well as of technological interest for noise characterization of future

spintronic devices, being closely related with magnetization reversal phenomena [153]. In

this context, the Barkhausen avalanche has attracted increased interest exhibiting

dynamical critical scaling behavior which is evidenced by a power-law distribution of the

Barkhausen jump size [154] as shown in Figure 6.4. Similar scaling behavior has been

observed in a wide variety of other physical phenomena such as lung inflations, mass

extinctions, micro fractures, and earthquakes [155]. The most interesting question in the


157

Figure 6.3 Series of local displacement having discrete and jerky jumps during

magnetization reversal revealing Barkhausen effect


158

Figure 6.4 The power law scaling behavior


159

research of critical scaling phenomenon is whether there exists a scaling exponent of the

power law distribution and whether these exponents are universal.

6.6 The Critical Scaling Behavior in 2D Ferromagnetic Films

Most of the experimental studies on Barkhausen avalanche have been carried out

on bulk samples employing a classical inductive technique which is difficult to apply to

thin films mainly due to the low signal intensity of the films [156]. This is the reason for

the very few experiments that have been performed, so far on 2D (two-dimensional)

ferromagnetic thin films. Co and MnAs thin films with uniaxial in-plane anisotropy have

been found to exhibit critical scaling behavior quite successfully [155]. The Barkhausen

avalanches at criticality were directly visualized using a novel instrument, named the

magneto-optic microscope magnetometer (MOMM). The scaling exponent in Co films

was found to be ≈ 1.33, independent of film thickness while the value of critical exponent

in MnAs films was found to vary from 1.32 to 1.04, depending on temperature.

6.7 The Critical Scaling Behavior in 2D γ′-Fe4N Thin Films

The main purpose of this study was to investigate the question whether there is any

universality of the critical exponent for the Barkhausen avalanches at criticality in the

epitaxial and non-epitaxial γ′-Fe4N thin films. For this purpose the epitaxial and non-

epitaxial this film samples were prepared by dc magnetron sputtering directly on

MgO(100) substrates. The detailed deposition conditions and preparation method has

been described in section 4.1. The crystal structure of the samples was confirmed by

XRD as shown in the Figure 4.3.

The MOMM system as described in detail in section 3.11 was utilized in order to

explore the Barkhausen criticality of the thin films. Since the MOKE is a surface

sensitive technique, therefore the MOMM is an ideal system to monitor magnetization

process in magnetic thin films. The Barkhausen avalanche was triggered by applying a

constant magnetic field to an initially saturated sample. The strength of an applied field

was about 99% of the coercivity field to eliminate the influence caused by the difference

in the field sweeping rate. Using MOMM, the Barkhausen avalanches were directly

visualized in real time and characterized from serial time resolved domain images.


160

Figures 6.5 – 6.8 show the series of four representative domain-evolution patterns

of γ′-Fe4N thin films observed on exactly the same area of the samples captured

successively using MOMM. The color code from red to blue represents a time lapse of

4 sec, as the color palette at the bottom of the figures indicates. Each time, the sample

was initially saturated, and a constant field H was then applied in the opposite direction.

The patterns shown in Figures 6.5 – 6.7 were obtained from the samples deposited at 450

˚C and in-situ annealed for 10-30 minutes, respectively, exhibiting epitaxial growth. The

Figure 6.4 (d) shows the patterns obtained from the sample deposited at 450 ˚C and in-

situ annealed for 40 minutes exhibiting non-epitaxial growth. It can clearly be noted that

magnetization reverses with a sequences of discrete and jerky jumps rather than a

continuous and smooth process. Therefore, the Barkhausen avalanches can directly be

witnessed in the evolution patterns of γ′-Fe4N films. Moreover, It can also be observed in

the domain evolution patterns among four pictures in any one set taken exactly on the

same area of the sample seem to display randomness with respect to interval, size and

location of the jump.


161

Figure 6.5 A Series of six domain images showing Barkhausen avalanche in the film

annealed for 10 min


annealed for 20 min


162


annealed for 30 min


annealed for 40 min


163

A direct visualization of the domain evolution by means of the MOMM enables

us to witness the randomness of magnetization reversal process at Barkhausen criticality.

Interestingly, it can be noted that a simple 180˚ flexible domain wall exists throughout

the Barkhausen avalanche process, in all the samples, irrespective of the structural nature

of the films related to their epitaxy.

To answer the fundamental question whether there is a simple law which governs

the randomness of the Barkhausen avalanches observed in the 2D iron nitride films, a

statistical analysis is performed by taking a large number of measurements. For this

purpose, more than 1000 repetitive experiments were carried out for each sample. The

Figures 6.9 - 6.12 demonstrate the plots of the distribution P(s) vs jump size „s‟ in the

log-log scale for the iron nitride samples deposited at 450 ˚C and in-situ annealed from

10 to 40 minutes. As can be seen in the figures, a power-law scaling distribution of the

jump size is found to exist for all the samples and fitted as P(s) s-τ with the critical

exponent τ = 1.02±0.07, 1.00±0.06, 1.04±0.06, and 1.03±0.05, for the films annealed for

10, 20, 30, and 40 minutes respectively. Our previous structural investigations on these

samples have already proved the epitaxial nature of the samples annealed for 10 to

30 minutes and non-epitaxial nature of the sample annealed for 40 minutes [50].

An interesting feature of the Barkhausen criticality observed in 2D thin films of

iron nitride is that the critical exponent is invariant with respect to the structural nature

related with the epitaxy of the iron nitride films as shown in the Figure 6.13. This

universality of the critical exponent seems to be very striking as both the epitaxial and

non-epitaxial samples of iron nitride have displayed the value of critical exponent in the

same range. Naturally, it is expected that epitaxial films being grown smoothly having

layer by layer growth texture have a smaller number of defects as compared to the non-

epitaxial films. The non-epitaxial films obviously have more imperfections at

polycrystalline grain boundaries. In this context, the experimental results obtained for

different types of films imply an invariance of the critical exponent irrespective of the

number of defects in the 2D iron nitride films. These findings are consistent with the

theoretical studies predicting that the variance of the number of defects does not affect

the critical exponent [157].


164

Figure 6.9 The distribution of Barkhausen avalanche size for the sample annealed

for 10 min.

0.01 0.1

1

10

100

P(s

) (a

.u.)

Jump Size, s (a.u.)


165


for 20 min.

0.1

1

10

100

P(s

) (a

.u.)

Jump Size, s (a.u.)


166


for 30 min.

0.01 0.1

1

10

100

P(s

) (a

.u.)

Jump Size, s (a.u.)


167


for 40 min.

0.01 0.1

1

10

100

P(s

) (a

.u.)

Jump Size, s (a.u.)


168

Figure 6.13 Effect of annealing time on critical exponent τ observed in 2D

iron nitride thin films

10 20 30 400.8

0.9

1.0

1.1

1.2

1.3

Exp

onent


Chapter 7 Conclusions

169

Conclusions

As the title of this thesis demands, an extensive effort has been made to explore

the ferromagnetic iron nitride system. For this purpose, γ′- and α″-phases in thin film

form have been deposited on different single crystal substrates. The cause of high

magnetic moment in this system is attributed to the unique structural properties of the

films. In this chapter, the conclusions drawn based on the present research have been

described.

At the first stage, we deposited γ′-Fe4N thin films on MgO(100) substrates at

various conditions of substrate temperature and annealing time. XRD of the samples

deposited at 450 ˚C and in-situ annealed for 10 – 40 minutes revealed the most favorable

conditions for the highly structured epitaxial growth of the films. The epitaxial texture

was improved as the annealing time was increased from 10 to 30 minutes. Annealing the

sample for 40 minutes destroyed this epitaxial nature which is due to excessive escape of

nitrogen due to longer annealing time. The rocking curve analysis, pole figure and phi

scan confirmed that 450 ˚C substrate temperature and in-situ annealing of 30 minutes are

the most suitable conditions for the single phase epitaxial growth of γ′-Fe4N thin films.

A regular nano-sized grain structure turned gradually in to smooth layered surface

as the temperature was increased from 200 ˚C to 450 ˚C. This gradual change in surface

morphology is caused by the increased surface mobility as the temperature is increased.

Furthermore, surface mobility is reduced at exceedingly higher temperatures, which

happened too in our case. At 500 ˚C, a reduced grain size was obvious.

The full width at half maximum and the RMS value of surface roughness were

observed to be highly correlated. With the decrease of annealing time from 10 to

30 minutes, both the values decreased and reached to 0.37˚ and 0.17 nm, respectively.

This trend is another indication of increased epitaxy of the films. At 40 minutes

annealing, both the RMS roughness and the FWHM increased sharply which is also

consistent with the XRD patterns of the films.

These structural characteristics of the γ′-Fe4N films were found to have a direct

impact on the magnetic properties of the films. The Ms was enhanced as the texture of the

films, in terms of FWHM and RMS roughness, improved by increasing annealing time.


170

The maximum Ms was observed for the sample annealed for 30 minutes, corresponding

to minimum values of FWHM and RMS roughness. The strong ferromagnetic behavior

of this sample was also obvious from the static and magnetic domain images. Bi-axial

cubic magnetic anisotropy was witnessed from the approximately perfect square

hysteresis loops. A dense TEM-microstructure and electron holographic images exhibited

strong exchange coupling among magnetic domains. When the temperature was

decreased below RT, a slight increase in magnetic moment was noticed which was

expected as moments tend to align perfectly at low temperatures. The metallic behavior

of the films was evident as the resistivity decreased by decreasing temperature below RT.

Until this stage, our findings depicted that the enhancement of magnetization was

strongly dependent on the epitaxial texture of γ′-Fe4N films. Therefore, to confirm this

observation, we have investigated the effect of lattice mismatch on Ms, by depositing γ′-

Fe4N thin films on MgO(100), SrTiO3(100), and LaAlO3(100) substrates having lattice

mismatches of 11%, 3%, and 0%, respectively. A pronounced increase in Ms was

observed with decrease in lattice mismatch and the film grown on LaAlO3(100) substrate

having 0% lattice mismatch exhibited maximum Ms of 1980±20 emu/cc, inferring that

the enhancement of Ms is directly related with minimizing the lattice mismatch and hence

on the epiatxy of γ′-Fe4N films.

The controversy about the giant magnetic moment of α″-Fe16N2 has always been a

key issue in the iron nitride system. Epitaxial thin films of α″-phase have been reported to

have the highest value of saturation magnetization. However, some research groups could

not reproduce those results. We have achieved optimum conditions for the single phase

epitaxial growth of α″-Fe16N2 by depositing these films on Si(100) substrates by varying

nitrogen partial pressure. All the films were deposited at 200 ˚C and in-situ annealed for

1 hour. Single phase epitaxial growth was achieved for a nitrogen partial pressure of

0.8 m Torr as observed by XRD. A high Ms value of 1800±20 emu/cc was observed for

this sample. However, this value is less than the previously reported value for this phase.

Since all the samples in this series were deposited at the same temperature and annealed

for the same period of time, the surface images of all the samples showed the same

morphology with a grain size ranging from 15-20 nm. The RMS value of surface

roughness was less for the epitaxially grown film as compared with the other non-

epitaxial films of the series.


171

The alloying addition of certain elements might be the solution to stabilize α″-

phase thermally. For the purpose, FeCoN, FePtN, and FeCrN alloy phases were prepared.

The at% concentration of alloying elements, as revealed by the AES spectra, in the iron

nitride matrix increased almost linearly by increasing the number of chips placed on the

iron target. For the FeCoN system, it is inferred that the alloying addition of Co in α′′-

Fe16N2 phase could be favorable as for as stability is concerned, with some compromise

on saturation magnetization. More systematic investigations are needed to explore FePtN

alloy phases to make this system compatible for applications that require high Ms with

relatively high coercivity. Ferromagnetic-antiferromagnetic coupling between Cr and Fe

might the cause for decreasing Ms with increase in Cr contents in FeCrN system.

The Barkhausen effect, a characteristic phenomenon of ferromagnetic materials

has also been investigated in 2D γ′-Fe4N thin films. The samples selected for this study

were composed of epitaxial and non-epitaxial thin films of iron nitride as described in

section 4.2. The discrete and jerky jumps known as “Barkhausen jumps” are the main

cause of noise heard in magnetic devices during magnetization reversal process. The

random distribution of jump size was found to obey the “power law scaling behavior”

with a critical exponent τ = 1.02±0.07, 1.00±0.06, 1.04±0.06, and 1.03±0.05, for the films

annealed for 10, 20, 30, and 40 minutes respectively.

The universality of the critical exponent found in the epitaxial and non-epitaxial

films showed that the critical exponent is invariant, irrespective of the number of defects

in 2D γ′-Fe4N films, as the numbers of defects in non-epitaxial films are generally more

as compared with epitaxial films. These findings are also consistent with the theoretical

results on the subject.

References

172

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