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2015

Optical characterization of ferromagnetic and multiferroic thin-Optical characterization of ferromagnetic and multiferroic thin-

film heterostructures film heterostructures

Xin Ma College of William & Mary - Arts & Sciences

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Recommended Citation Recommended Citation Ma, Xin, "Optical characterization of ferromagnetic and multiferroic thin-film heterostructures" (2015). Dissertations, Theses, and Masters Projects. Paper 1539623372. https://dx.doi.org/doi:10.21220/s2-66yg-6612

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Optical Characterization of Ferromagnetic and Multiferroic Thin-Film Heterostructures

Xin Ma

Lingbi, Suzhou, Anhui, P. R. China

Bachelor of Science, the University of Science and Technologyof China, 2009

A Dissertation presented to the Graduate Faculty of the College of William and Mary in Candidacy for the Degree of

Doctor of Philosophy

Department of Applied Science

The College of William and Mary January, 2015

APPROVAL PAGE

This Dissertation is submitted in partial fulfillment of the requirements for the degree of

Doctor of Philosophy

10 I ' I o-Xin Ma

Approved by the Committee, November, 2014

Committee Cnair Professor Gunter Luepke, Applied Science

The College of William and Mary

Professor Michael Kelley, Applied Science The College of William and Mary

Professor Mark Hinders, Applied Science The College of William and Mary

Professor Mumtaz Qazilbash, Physics The College of William and Mary

ABSTRACT

This thesis presents optical characterization of the static and dynamic magnetic interactions in ferromagnetic and multiferroic heterostructures with time-resolved and interface-specific optical techniques. The focus of the thesis is on elucidating the underlying physics of key physical parameters and novel approaches, crucial to the performance of magnetic recording and spintronic devices.

First, time-resolved magneto-optical Kerr effect (TRMOKE) is applied to investigate the spin dynamics in L1o ordered FePt thin films, where perpendicular magnetic anisotropy Ku and intrinsic Gilbert damping ar0 are determined. Furthermore, the quadratic dependence of Ku and a0 on spin- orbit coupling strength f is demonstrated, where f is continuously controlled through chemical substitution of Pt with Pd element. In addition, a linear correlation between cro and electron scattering rate 1 /re is experimentally observed through modulating the anti-site disorder c in the L1 0 ordered structure. The results elucidate the basic physics of magnetic anisotropy and Gilbert damping, and facilitate the design and fabrication of new magnetic alloys with large perpendicular magnetic anisotropy and tailored damping properties.

Second, ultrafast excitation of coherent spin precession is demonstrated in Fe/CoO heterostructures and Lao.6 7Cao.3 3Mn0 3 thin films using TRMOKE technique. In the Fe/CoO thin films, Instant non-thermal ferromagnet (FM) - antiferromagnet (AFM) exchange torque on Fe magnetization through ultrafast photo-excited charge transfer possesses in the CoO layer is experimentally demonstrated at room temperature. The efficiency of spin precession excitation is significantly higher and the recovery is notably faster than the demagnetization procedure. In the Lao.67Cao.33Mn03 thin films, pronounced spin precessions are observed in a geometry with negligible canting of the magnetization, indicating that the transient exchange field is generated by the emergent AFM interactions due to charge transfer and modification of the kinetic energy of eg electrons under optical excitation. The results will help promoting the development of novel device concepts for ultrafast spin manipulation.

Last, the interfacial spin state of the multiferroic heterostructure PbZro.5 2Tio.4 8O3/Lao.6 7Sro.3 3MnO3 and its dependence on ferroelectric polarization is investigated with interface specific magnetization induced second harmonic generation (MSHG). The spin alignment of Mn ions in the first unit cell layer at the heterointerface can be tuned from FM to AFM exchange coupled, while the bulk magnetization remains unchanged as probed with MOKE. The discovery provides new insights into the basic physics of interfacial magneto-electric (ME) coupling.

TABLE OF CONTENTS

Acknowledgements iii

Dedication iv

List of Tables v

List of Figures vi

Chapter 1. Introduction

1.1 Magnetic recording and spintronics 1

1.2 Magnetization switching and dynamics 3

1.3 Interface magnetization 6

1.4 Scope of dissertation 9

Chapter 2. Experimental techniques

2.1 Magneto-optical Kerr effect 12

2.2 Time-resolved Magneto-optical Kerr effect 17

2.3 Magnetization induced second harmonicgeneration 26

Chapter 3. Time resolved studies of spin dynamics in FePt thin flims

3.1 Introduction 34

3.2 Samples and characterizations 36

3.3 Magnetic anisotropy and damping inFePtxPd(i-x) thin films 43

3.4 Magnetic anisotropy and damping in FePtthinfilms with different chemical order 55

Chapter 4. Time resolved spin precession in Fe/CoO heterostructure

4.1 Introduction 67

i

4.2 Samples and characterizations 69

4.3 Magnetic anisotropy 72

4.4 Ultrafast spin exchange torque via photo­excited charge transfer processes 82

Chapter 5. Spin precession excitation in manganite thin films

5.1 Introduction 100

5.2 Samples and experiments 102

5.3 Coherent Spin Precession via Photoinduced Anti-ferromagnetic Interactions in Lao.6 7Cao.3 3Mn0 3 103

Chapter 6 . Magneto-electric coupling at PbZr0 5 2Ti 0 .4 8 O3 /Lao 6 7Sr0.3 3MnO3 interface

6.1 Introduction 117

6.2 Samples and experiments 119

6.3 Charge control of antiferromagnetism atPbZr0 52Ti 0 4 8 O3/ La0.67Sro.33Mn03 interface 124

Chapter 7. Conclusions 136

Bibliography 139

Vita 156

ii

ACKNOWLEDGEMENTS

I am truly grateful to everyone who has directly or indirectly supported and helped me during the completion of this dissertation. First and foremost, I would like to thank my advisor, Professor Gunter Luepke, for his guidance, encouragement and mentoring throughout my doctoral study at the College of William and Mary. He has always provided enthusiasm and perspective for the research process. This dissertation would not have been possible without his support and guidance.

I am particularly grateful to Professor Haibin Zhao for guiding my experiment in the lab and fruitful scientific discussions. I truly appreciate his sharing of academic experience with me as a friend.

I owe special thank to Professor Shiming Zhou and his group members - Pan He and Li Ma at Tongji University in Shanghai, China. I am indebted to them for providing us state-of-art FePt thin films, structure characterizations and inspired visions.

I truly appreciate Professor Yizheng Wu and his group members- Qian Li and Jie Zhu at Fudan University in Shanghai for their support and providing us high-quality Fe/CoO thin films.

I am similarly grateful to Professor Ram Katiyar and Dr. Ashok Kumar for providing us high-quality multiferroic heterostructures and characterization results. I also appreciate Professor James Scott for helpful suggestions.

I am indebted to Professor Qi Li from Penn State University, for her generosity in providing us high-quality manganite films and sharing academic experience.

My sincere gratefulness goes to my lab mates: Fan Fang, Wei Zheng, Haowei Zhai. I appreciate all their friendships and their collective encouragement to finish this dissertation. I have enjoyed every moment that we have worked together.

I would also like to express my appreciation to my doctoral dissertation committee members: Professor Michael Kelley and Professor Mark Hinders of Applied Science, Professor Mumtaz Qazilbash of Physics for their comments and suggestions during the course of completing this dissertation.

Finally, I wish to express my sincere gratitude to my parents - Renmi Ma and Ling Huang, and my wife - Yufei Ding. Their constant love, encouragement and support have always been the source of my strength and motivation. Anything I have been able to accomplish is a tribute to them.

This dissertation is dedicated to my family.

LIST OF TABLES

2.1 Fresnel coefficients, where = cosO, 6 the light incidentangle, a2 = J l - sin20 /n 2, QP,QT,QLthe polar, transverse,and longitudinal projections of the Voigt vector Q, respectively.

2.2 Nonzero tensor elements of the nonlinear susceptibility tensor * (2) for the (0 0 1) cubic surface, the magnetization parallel to the x-axis (longitudinal configuration), y-axis (transversal configuration) and parallel to the z-axis (polar configuration). * (+) and denote the even and odd elements with magnetization. Other non-zero elements with mixed polarization of s and p are not included. For simplicity of notation the tensor components are indicated by their indices only

3.1 Landau g factors and Ku derived from TRMOKEmeasurements for different doping level x at 20K and 200K.

LIST OF FIGURES

1.1 Schematics of magnetic recording arrays with traditional writing mechanism.

1.2 Schematic of precessing magnetization.

1.3 Relative orientations of the atomic moments in the AFM and FM are shown schematically, illustrating the lateral offset in the magnetization curve.

2.1 Schematics of longitudinal, transverse and polar MOKE probing geometry, where POS is the plane of sample (x-y plane), and POI is the plane of incident (y-z plane).

2.2 Geometry of TRMOKE measurements.

2.3 Stages of the excitation process: (a) M aligns along the equilibrium direction H eff \ (b) new equilibrium directionH'eff when pump pulse “heat up” the sample; (c) Mprecesses around H'eff (d) heat diffused and Mprocesses around H eff.

2.4 TRMOKE result from L10 FePt thin film with growth temperature 550 °C at magnetic field 3.5T.

2.5 Schematic of TRMOKE measurement geometry for L1 o FePt thin films.

2.6 Coordinate system used for calculating the magnetization precession frequency in the L10 FePt thin films.

2.7 Schematic of TRMOKE experimental set-up.

2.8 Schematic of MSHG and MOKE measurements on PZT/LSMO thin film.

2.9 MSHG hysteresis loop from PZT/LSMO thin films indicating l(-M) and l(+M) used to determine the magnetic contrast A from Eq. (2.17).

2.10 The schematic set-up of MSHG measurement.

3.1 Schematics of atomic structures for the two series ofFePt thin films.

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3.2 Structure characterization results of FePtxPd(i.X) thin films(a) and FePt thin films with different Tg (b) by X-ray Diffraction measurements.

3.3 Anti-site disorder percentage c in Z_10 ordered FePt thin films as a function of growth temperature Tg.

3.4 Static magnetic hysteresis loops measured by vibrating sample magnetometer (VSM) (a)-(c) for FePtxPd(i-x) thin films. The black (red) curves represent out-of-plane (in­plane) results. VSM results for FePt thin films with different Tg (d).

3.5 TRMOKE data with different doping level of FePtxPd(l-x) under magnetic field H=5T. The solid lines are fitted curves.

3.6 TRMOKE results of FePt0.82Pd0.18 (a) and FePt0.5Pd0.5 (b) under different magnetic fields H. The solid lines are fitted curves.

3.7 The dependence of precession frequency (a) and lifetime(b) on external magnetic field H for typical samples. The solid lines refer to fitted results.

3.8 PMA (a) and ao (b) as a function of doping level x of FePtxPd(i-x)

3.9 Spin-orbit coupling strength f (a), bandwidth W (b), spin moment ms on Fe or Pt/Pd site (c) and measured c/a ratio (d) as a function of doping level x of FePtxPd(i.X).

3.10 Measured and calculated ao (a) and PMA (b) as a function of SOC strength. The solid line is a quadratic fit obtained from first principle calculations.

3.11 TRMOKE data from 22-nm thick FePt films, grown at different Tg with applied magnetic field H=6.5T (a), and with Tg =680 °C measured at different H (b). The solid lines refer to fitted results.

3.12 The dependence of spin precession frequency f (c) and effective Gilbert damping aeff (b) on H obtained from Eqs.3.3 and 3.5 for FePt thin films with different x. The solid lines are fitted results.

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3.13 PMA Ku (a) and Landau g factor (b) as a function of anti­site defect concentration c for FePt films with thicknesses of 17 nm and 22 nm.

3.14 ao (a), and ao/g (b) as a function of anti-site defect concentration c for FePt films with thicknesses of 17 nm and 22 nm. The solid line refers to least square linear fit.

3.15 The dependence of f and 1/r on temperature T with applied field H =5T (a), and as a function of H at 20 K (b). The solid lines refer to fitted results.

4.1 Longitudinal hysteresis loops for magnetic field along uniaxial easy axis [100] and hard axis [010] at 82 K and 330 K.

4.2 Temperature dependence of coercivity (Hc) and exchange biasing field (He) obtained from easy axis loops.

4.3 TRMOKE results of Fe/CoO thin films with different magnetic field H along the Fe [110] direction (in-plane hard axis) at RT.

4.4 Spin precession frequency f as a function of magnetic field H.

4.5 Temperature dependent TRMOKE results with magnetic field H= 3 kOe.

4.6 Temperature dependent TRMOKE results with magnetic field H= 1 kOe.

4.7 Spin precession frequency fas a function of temperature.

4.8 (p as a function of temperature (green=3 kOe, blue=1 kOe).

4.9 9 as a function of temperature (green=3 kOe, blue=1 kOe).

4.10 Uniaxial magnetic anisotropy field KJMs as a function of temperature.

4.11 Two pump strategies to optically excite the spin precession, where the black arrows represent the spins of magnetic atoms and ions. The optical charge transfer transition in CoO is depicted.

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4.12

4.13

4.14

4.15

4.16

4.17

4.18

4.19

5.1

5.2

5.3

TRMOKE results from Fe/MgO (black squares) and Fe/CoO (red circles) with pump pulse intensity 3.1 mJ/cm2, and Fe/CoO (blue triangles) with pump intensity 0.16 mJ/cm2.

Measured and simulated spin precession amplitude A as a function of T.

The spin precession amplitude A is plotted as a function of H at different T, where solid lines represent simulated results.

Measured Kerr signal change with 0.8 ps step size (black squares), simulated polar magnetization component (red curve) and assumed modulation of Ku (blue curves) as a function of t.

Measured Kerr signal change with 4 ps step size (black squares), simulated polar magnetization component (red curve) and assumed modulation of Ku (blue curves) as a function of t.

Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with 77 = 2 ps.

Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with 7> = 15 ps.

Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with rr = 1 0 0 ps.

Transient reflectivity AR (a), and transient Kerr rotation A0 of LCMO film with magnetic field H (0.5 T) applied perpendicular to the sample plane (b), with H applied along the in-plane easy axis (c), and with zero field (d) at 20 K; Fourier transform (e) of Fig. 5.1(d).

Transient Kerr rotations of Fe film with magnetic field applied along the in-plane hard axis (a) and easy axis (b) at RT.

Three-temperature model in ferromagnetic metals and half-metals.

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5.4 Amplitude of precession of 60 nm and 100-nm LCMO films with field applied at 45 degree to the easy axis and along the easy axis at 20 K (a), and Fe film with field applied at 55 degree to the easy axis and of 10 degree to the easy axis at RT (b).

5.5 Transient exchange field (a) and transient anisotropy field (b) in a state of canted magnetization (magnetic field at 5 degree to the easy axis) of 60-nm thick LCMO film. The dashed line is a guide to the eye of the data.

5.6 Transient Kerr rotations of LCMO film with applied field H=0.5T at various temperatures. The inset shows the uniform precession frequency as a function of the temperature.

6.1 Schematic of MSHG and MOKE measurements (a), and depiction of PZT/LSMO interface on atomic scale (b). MSHG probing plane is marked with purple color.

6.2 XRD results of PZT/LSMO heterostructure (a), RHEED patterns of LAO substrate surface (b), LSMO surface before deposition of PZT layer on top (c) and PZT surface of PZT/LSMO/LAO heterostructure (d).

6.3 Interface (a) and Bulk (b) magnetic hysteresis loops from PZT/LSMO heterostructure for different external gate voltages Ug measured with MSHG and MOKE techniques at 78 K, respectively.

6.4 MSHG hysteresis loop with Ug = +10V indicating l(-M) and l(+M) used to determine the magnetic contrast A from Eq. (6.1) (a), and MSHG magnetic contrast A as a function of Ug (b). Increasing (decreasing) gate voltages, Ug up (Ug down), are labeled in red (black). The direction of positive Ug is defined in the inset.

6.5 Ferroelectric polarization (P) in PZT as a function of Ug (a). The direction of positive P is defined as pointing towards LSMO thin film. MSHG magnetic contrast A as a function of P of PZT (b). The dashed line in Fig. 6.5(b) is the least square fit for the data points.

6.6 Strain state of LSMO interface as a function of gate voltage (a). MSHG magnetic contrast A as a function of strain (5) of PZT (b).

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6.7 Schematic model of ions displacement and screen charge accumulation under opposite polarization state.

6.8 P in S out MSHG results under different polarization states.

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Chapter 1

Introduction

1.1 Magnetic recording and spintronics

For the past few decades, magnetic recording and spintronics have

been the major technologies for information storage and processing [1-2]. The

magnetic devices have been widely applied in the hard disk read heads,

memories and sensors, where the binary information is represented by the

magnetization of magnetic elements, instead of electrons or holes in

traditional semiconductor electronics [3]. Nowadays, the large amount of

digital information creates ever-increasing demands for much higher storage

densities, faster and more energy-efficient processing operations. Therefore,

it attracts abundant research interest to investigate the key physical

parameters and mechanisms for the development of device performance.

The storage density is closely related to the magnetic anisotropy of the

magnetic element [4-6]. Each piece of magnetic recording media carries a bit

of information through controlled magnetization directions, while the element

size is limited by the thermal fluctuations of the magnetization directions [7, 8].

Magnetic anisotropy provides the energy barrier between these states,

allowing them to withstand the loss of stored information due to thermal

fluctuations. Therefore, the search for material structures with larger magnetic

anisotropies promotes the increase of storage densities, as shown in Fig 1.1.

l

In order to improve the speed and energy consumption of the

processing operations, novel approaches to effectively manipulate the

magnetization in the magnetic devices are investigated. For example,

optically generated effective magnetic field pulses and the spin-transfer

torque (STT) effect from spin-polarized carrier transport are recently applied

to switch the magnetization in spin valves and magnetic tunneling junctions

(MTJs) [9-12]. These techniques can be very strong, interact locally [13], and

thus are favorable for the magnetization manipulation in high-density

magnetic devices. With respect to energy consumption, it is more energy-

favorable to generate electric fields instead of currents in highly integrated

devices. As a result, the electric control of magnetization offers new pathways

to improve device performance. [14] Multiferroic materials and

heterostructures are of particular interest for using electric fields to control the

magnetization via the ME coupling effect [15].

Figure 1.1 Schematics of magnetic recording arrays with traditional writing

mechanism [16].

*Ring* writing element

Longitudal Recording (standard)

] Recording layer

'Monopole’ writing element

Perpendicular Recording

La Recording Layer | Additional Layer

2

1.2 Magnetization switching and dynamics

1.2.1 Magnetization switching process

In the demand for faster operation, the dynamic properties of

magnetism in magnetic devices and controlled magnetization switching

become increasingly important. The magnetization switching process is well

modeled with the Landau-Lifshitz-Gilbert (LLG) equation [17],

d M m m > • M C^M / jt jk \— = - YM x H eff + a - x — (1.1)d t 1 M d t ' ’

where the first term on the right indicates that the magnetization M does not

go directly to the desired alignment along H eff, but instead, circles around

H ef f ■ The second term describes the damping of the circling precession

towards H eff with the Gilbert damping parameter a, where larger a

represents faster relaxation and dissipation of magnetic energy into the

thermal bath. Therefore, the speed and damping of coherent spin precession

play important roles in the operation speed of magnetic devices.r

i

Figure 1.2 Schematic of precessing magnetization.

3

1.2.2 Damping mechanisms

Well-controlled magnetic relaxation is required for designing magnetic

devices with promising performance. For most magnetic transition metals, a is

small; so that M precesses back and forth several times after switching or

small perturbation, which takes several nanoseconds to relax the precession.

This actually puts a bottleneck on the switching speed. Therefore in many

applications requiring high speed or high data rates, the stronger damping

behavior is desirable. On the other hand, larger cr indicates stronger coupling

between the magnetic system and the thermal bath. This would allow the

thermal bath to drive fluctuations of the magnetization, and cause

magnetization noise in small magnetic sensors. For these applications, the

noise limit could be improved with reduced damping. Thus in all, some

sophisticated methods to control the magnetic relaxation behavior are desired.

The Gilbert damping parameter a contains both extrinsic and intrinsic

components. The extrinsic Gilbert damping is due to nonlocal spin relaxation,

such as spin pumping and two-magnon scattering. This extrinsic component

can be tuned by artificial substrates, specially designed buffer and coverage

layers. On the other hand, the intrinsic Gilbert damping parameter a0 is due to

a local magnetic energy transfer effect and describes the energy flow rate

from spin to electronic orbital and phonon degrees of freedom through

electron scattering.

4

1.2.3 Excitation of Magnetization switching

The conventional method to switch the magnetization alignment is

magnetizing the storage cell with an external magnetic field. As mentioned

above, the increasing storage density requires larger magnetic anisotropy,

which provides a significant energy barrier for switching the magnetization

between the spin-up and spin-down states. The switching actually requires

the application of stronger magnetic field, leading to larger energy

consumption. Recently, a breakthrough of this limitation is the utilization of

optical demagnetization or spin-polarized current to overcome the large

magnetic energy barrier. For instance, effective magnetic field pulses can be

generated by the optical modulation of magnetic anisotropies with laser

pulses, which trigger the magnetization switching. Another example is the

spin-transfer torque (STT) effect generated by optically or electrically

triggered spin-polarized carrier transport. The STT effect is due to the

interaction between itinerant electrons in FM thin films and the magnetization

[13], and can be used to switch the magnetization.

The critical density of the spin polarized current for STT switching is

proportional to aKuMs [9-12], where a is the Gilbert damping parameter, Ku

the magnetic anisotropy and Ms the saturated magnetic moment. Considering

the storage density, the speed and energy consumption of magnetization

switching in STT devices, a smaller or along with large value of Ku is essential

for promising performance.

5

1.3 Interface magnetization

1.3.1 Magneto-electric coupling

The switching of magnetic state is accomplished by applying an

electric current, either for the generation of a magnetic field or for the STT

effect. The large current densities required also lead to significant energy loss

from heating. In contrast, it is more convenient and energy-favorable to

generate an electric field in high-integrated devices. Thus in the push for low-

energy-consumption logic devices, strong ME coupling seems to offer a new

mechanism to break through those limitations and improve the device

performance; as such, it has recently drawn increasing research interest.

In the search for stronger ME coupling, multiferroic structures are

promising candidates, where the materials exhibit both FM and ferroelectric

(FE) order. These materials are usually transition metal oxides, where the

competition between different interactions in such strong-correlated electron

systems results in a complex phase diagram, and provides a platform to

manipulate the coupling between different order parameters. For example in

some manganites [18], there are some phases with the coexistence of FM

and FE order. However, such phases usually exist at low temperature, and

the ME coupling is fairly weak, which prevents them from real-life applications.

To overcome these disadvantages, an alternative but challenging approach is

the ME coupling across the interfaces of multiferroic heterostructures, which

contain room-temperature FE and FM order. Moreover, the electric control of

6

interface magnetization can be effectively incorporated into devices where the

performance depends crucially on the interface magnetization, such as MTJs.

FM

A FM

(ii)

(iii)

(Hi) (iv)

(iv )

-> H

Figure 1.3 Relative orientations of the atomic moments in the AFM and FM are

shown schematically, illustrating the lateral offset in the magnetization curve.

7

1.3.2 Exchange bias

Exchange bias occurs at the interface between an AFM layer and a FM

layer, where the AFM order biases the magnetization of the softer FM layer.

Exchange bias can be generated by field-cooling the AFM/FM structure with a

magnetic field applied through the Neel temperature of the AFM layer (the

temperature at which AFM order sets in). The very strong exchange coupling

between the FM and AFM layers tends to pin the FM magnetization in a

specific direction, which can also be interpreted as the generation of uni­

directional magnetic anisotropy. This results in an offset of the hysteresis loop,

so that it is no longer centered at zero applied field, but shifted by an amount

corresponding to the exchange bias field He, as illustrated in Fig. 1.3. The

exchange bias increases the magnitude of the applied magnetic field needed

to reverse the FM magnetization from the normal coercive field Hc to Hc+ He.

Exchange bias is of great technological importance in tailoring the

operating characteristics of most magnetic devices, including hard disk read

heads, magnetic memory and magnetic sensors [19]. Moreover, electric

control of exchange bias offers a novel way for multiferroic devices. For

example, some material systems exhibit strong coupling between AFM order

and FE order at room temperature, such as BiFe0 3 . Combination of BiFeC>3

layer with FM-ordered LaSrMnC>3 leads to indirect ME coupling FE-AFM-FM,

where the exchange bias between BiFe0 3 and LaSrMnC>3 layers can be well

controlled by electric field.

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1.4 Scope of Dissertation

Fundamental understandings of magnetic properties and spin

dynamics in ferromagnetic metal alloy thin films, ferromagnetic metal /anti­

ferromagnetic oxide heterostructure, and multiferroic complex oxide

heterojunction are the focus of this dissertation:

Ferromagnetic metal alloy thin films - spin dynamics in L1o ordered

FePt thin films is investigated with TRMOKE technique. The perpendicular

magnetic anisotropy Ku and intrinsic Gilbert damping ao of the material

systems are determined from the spin dynamics study. On one hand, the

quadratic dependence of Ku and ao on spin-orbit coupling strength f is

rigorously observed for the first time, where f is continuously controlled

through chemical substitution of Pt with Pd element. On the other hand, a

linear correlation between ao and electron scattering rate 1/re is

experimentally observed through modulating the anti-site disorder c in the L10

ordered FePt thin films.

Ferromagnetic metal / anti-ferromagnetic oxide heterostructure - the

impact of FM-AFM exchange coupling on the spin precession excitation is

investigated in Fe/CoO heterostructure thin films with TRMOKE technique.

Instant non-thermal FM-AFM exchange torque on FM magnetization through

ultrafast photo-excited charge transfer processes in the AFM layer is

experimentally demonstrated at room temperature. The strong exchange

torque excites pronounced coherent spin precession with negligible

9

demagnetization in a Fe/CoO heterostructure, thus avoiding excessive

heating by laser pulses.

Transition metal oxide thin film - the coupling between magnetic and

other properties are investigated in Lao.6 7Cao3 3 Mn0 3, due to its complex

magnetic phase diagram. For the first time, pronounced coherent spin

precession is observed in a geometry with negligible canting of the

magnetization in ferromagnetic Lao6 7Ca0.3 3Mn0 3 thin films using TRMOKE

technique. This indicates that the transient exchange field is generated by the

emergent AFM interactions due to charge transfer and modification of the

kinetic energy of eg electrons under optical excitation.

Multiferroic complex oxide heteroiunction - the interfacial spin state of

the multiferroic heterostructure PbZr0.52Ti0.48O3/La067Sr0.33MnO3 and its

dependence on ferroelectric polarization is investigated with interface-specific

MSHG technique. The spin alignment of Mn ions in the first unit cell layer at

the heterointerface can be tuned from FM to AFM exchange coupled, while

the bulk magnetization remains unchanged probed with MOKE.

10

Chapter 2

Experimental techniques

This chapter presents three different optical characterization

techniques which have been employed to study magnetic and electronic

properties: MOKE, TRMOKE and MSHG.

Section 2.1 describes the MOKE technique and the optical geometries

for measuring magnetization vector components. The principles and

experimental setup of TRMOKE technique - an all-optical technique to study

spin dynamics - is described in section 2.2. The magnetization precession

and its mathematical derivation are also discussed in this section. Section 2.3

presents an interface-specific magnetization probe technique - MSHG,

including a detailed analysis of nonlinear susceptibilities in PbZro 5 2Tio 4 8 O3

/La0 6 7 Sro.3 3Mn0 3 and corresponding experimental consideration for probing

interface magnetization.

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2.1 Magneto-optical Kerr effect

In general, the polarization state of light reflected from a magnetized

material will change depending on the state of magnetization, resulting in a

modified ellipticity [20]. If the incident light is linearly polarized, this process is

known as MOKE, which was first discovered by John Kerr in 1877 [21]. The

origin of this effect is presently described in the context of either microscopic

quantum theory [22] or macroscopic dielectric theory [23],

Microscopically, this effect originates from spin-orbit coupling between

light and electron spin. A potential gradient \7V is generated by light, which

couples local spins,5, with momentum,P, adding an interaction term (W x

P) • S to the Hamiltonian [22], In the macroscopic picture, such interaction

gives rise to off-diagonal components of the dielectric tensor F, which are

determined by the state of magnetization in the magnetic medium:

1 iQz - iQ y - iQ z 1 iQxiQy iQx 1

(2 .1)

where e is the average dielectric constant, and Q = ( Qx, Qy, Qz) is the Voigt

vector, which is proportional to the magnetization [24].

From Maxwell’s equations, the normal modes of light propagation can

be categorized into left-circular polarized light with index of refraction nL =

n ( l - i <?•£), and right-circular polarized light with refraction index nR =

n( 1 +^Q • k), where n = V i is the average refraction index and k is the unit

12

vector of the light propagation direction. A linear polarized light can be

considered as a superposition of left-circular and right-circular polarized light

with equivalent amplitude. The difference between the real parts of nL and nR

in the medium will lead to a phase shift between the two modes and result in

a polarization rotation of the light propagating through the medium.

Meanwhile, the difference between imaginary part of nL and nR will cause the

different absorption rates for the two normal modes, affecting the ellipticity of

the light.

In the Kerr effect, a convenient decomposition scheme is that the

polarization is a combination of p- and s-polarized light, which define the

polarization of light parallel and perpendicular to the incident plane,

respectively. Therefore, MOKE can be represented by:

where Eg1 and £ jnare the p- and s-polarized components of incoming light,

Eput and E$ut are corresponding components for outgoing light, and the

matrix r is the scattering matrix with Fresnel coefficients.

MOKE can be categorized into longitudinal, transverse, and polar

MOKE, according to the geometry of magnetization (M) with respect to the

incident plane of light and film plane (Fig. 2.1): in the longitudinal geometry

the magnetization component lies in both the light incident plane and film

plane; in the transverse geometry the magnetization component is in the film

(2.2)

13

plane but perpendicular to the light incident plane; for the polar geometry, the

magnetization component is perpendicular to the film plane.

For different MOKE geometries, the Fresnel coefficients are distinct, as

summarized in Table. 2. 1. As a result, the longitudinal and polar components

of M change the polarization state of the light upon reflection. In contrast, the

transverse component of M only changes the intensity of the reflected beam,

and no ellipticity occurs. Therefore, different detection schemes are required

to measure different magnetization components. For example, for p-polarized

incoming light the change of s-polarization is sensitive to the longitudinal and

polar components of M , while one has to probe the change of p- polarization

to measure the transverse component of Af.

For the MOKE measurements, a beam of light first passes through a

linear polarizer and is then incident on the sample; after reflection from

sample, the beam passes through an analyzing polarizer (analyzer) and is

measured by a photodetector. In all my measurements, the cross-polarized

combination of polarizers is used to detect the longitudinal and polar

components of M. In other words, the analyzer angle is set close to 90° with

respect to the incident beam polarization.

14

polar

Figure 2.1 Schematics of longitudinal, transverse and polar MOKE probing

geometry, where POS is the plane of sample (x-y plane), and POI is the plane of

incidence (y-z plane).

15

Table 2.1: Fresnel coefficients, where ax = cos0, 6 is the light incident angle,

a 2 = yjl - sin26 / n 2, QP, QT, QLare the polar, transverse, and longitudinal projections

of the Voigt vector Q, respectively.

Component vv YpS — r sp ss

j-i,

Longitudinal

* ie w + * * i. ' '* wt 1 ,A' ■ . * a1-n a 2

n a 1- a 2n a x+ a 2

•. jf(naa+ a z) (n a 2+ a i) a t+ n a 2

Qf,na1tan02 «1-n a2i(na1+a2)(na2+a1) ax+na2

-WW...

i ;t ȣ} pry* ' ^

l^ C ^ /o ^ a i—thQ^itanBz ■ *

T ~~ . /.'.w*? ' <• I/ “ • I ;X ; i l

ax-n a 2ax+na2

16

2.2 Time- resolved Magneto-optical Kerr effect

TRMOKE is an all-optical method to investigate the spin dynamics in

the time domain [25], which utilizes the pump-probe technique. The pump

beam is used to excite the magnetization precession, while the probe beam

measures the decay of excited magnetization state with MOKE technique. By

controlling the delay between pump and probe laser pulses, a real-time plot of

magnetization precession is obtained.

_ PumpProbe TRMOKE

Substrate

Figure 2.2 Geometry of TRM OKE measurements.

17

2.2.1 Excitation of magnetization precession

Figure 2.3 depicts the stages of the excitation process by pump laser

pulses. Before the pump pulse reaches the sample t< to , the magnetization Af

lies along the equilibrium direction H eff , determined by the magnetic

anisotropy field, demagnetization field, exchange coupling, and external

magnetic field H. Here we assume the easy axis of magnetic anisotropy is

perpendicular to the film plane. When the pump pulse impinges on the

sample at t= to , the magnetic state of the sample is modified, which can affect

the magnetic anisotropy, exchange coupling, and/or the demagnetization field.

As a result, the equilibrium direction of M a t this stage is changed

which induces a torque on M to align it along the new direction. Within a short

time region fo^fo+IO ps, M precesses around the new direction. After the

original magnetic state of the sample is restored, t> fo+ 1 0 ps, the precession

continues around the equilibrium direction due to the initial displacement of M.

Figure 2.4 displays the TRMOKE result from L1o FePt thin films. The

excitation of magnetization precession is as depicted in Fig. 2.3. In the

TRMOKE measurement, we utilize the MOKE geometry sensitive to polar

component (p-in, s-out). Therefore, Figure 2.4 can also be interpreted as a

real- time plot of Mz as a function of delay time. The t m is the time for

demagnetization process, corresponding to Fig. 2.3 (b). During this process,

the electron temperature Te is first raised by photon-induced excitations (spin-

conserved), and then relaxed through the phonon assisted spin-flip scattering.

18

Meanwhile, the spin temperature Ts increases, resulting in a reduction of

magnetization. The te is the time for the recovery process, depicted in Fig.

2.3 (c). Afterwards, the magnetization starts to precess around the equilibrium

direction as depicted in Fig. 2.3 (d).

t> to + 1 0 psto < t< to + 1 0 ps

Figure 2.3 Stages of the excitation process: (a) M aligns along the equilibrium

direction H ef f \ (b) new direction H'eff when pump pulse “heat up” the sample; (c) M

precesses around H'e ff (d) heat diffused and M precesses around H e ff .

19

O)

0 20 40D elay tim e (ps)

Figure 2 .4 TRM OKE result from /_10 FePt thin film with growth temperature 550 °C at

magnetic field 3.5T.

2.2.2 Precession frequency and effective Gilbert damping

The precession motion of Af back to the equilibrium direction can be

phenomenologically modeled with LLG equation [17]:

where y = ggB/h is the gyromagnetic ratio, and a is the dimensionless Gilbert

damping constant. The effective field H eff is the sum of external magnetic

20

field, demagnetization field, anisotropy field, dipolar field, and exchange field.

The first term on the right describes the torque on the magnetic moment

produced b y H eff, which drives a precession motion around the direction of

the equilibrium direction H e/ / . The second term on the right describes the

damping of precession due to the exchange of energy and angular

momentum with environment, until the magnetization finally aligns parallel to

H eff . Figure 2.5 depicts an example of magnetization precession.

Figure 2 .5 Schematic of TRM OKE measurement geometry for L10 FePt thin films.

For small angle precession, Eq. (2.3) can be solved analytically with

the following steps: ( 1 ) finding the equilibrium orientation of magnetization

determined by H eff\ (2 ) deriving the equations of motion by describing the

torque on spins by the effective field; (3) solving the linearized equations of

motion, to obtain the frequencies and Gilbert damping of the spin wave

modes.

21

For example in L1o ordered FePt material system, the magnetic free

energy F is the sum of Zeeman energy, demagnetization and anisotropies:

F = —HMS cos(0 — (p) + 2ttM^ c o s 2 0 + Kusin20 (2.4)

where H is the magnitude of applied magnetic field, Ms is the saturated

magnetization, B and q> are the angles between M and [0 0 1 ] direction, and H

and [0 0 1 ] direction, respectively. Ku is the perpendicular magnetic anisotropy.

For convenience, the anisotropy field Ha = 2Ku/M s and demagnetization field

Hd = 4tcMs . The equilibrium direction for M can be derived by taking the

derivative of F over 0:

n _ s in (20)Ha- H d 2*s in ( (p -6 ) ' ' ’

i L

Z ’ //[001]z / / / I I

X ®

Figure 2 .6 Coordinate system used for calculating the magnetization precession

frequency in the L10 FePt thin films.

22

The magnetization precession motion can be written as:

d m x •* , d m y d m z—~ x + ——y + —d t d t d t

-»X y-*z

y mx my Ms + mzHx Hy Hz

(2.6)

where Hx,Hy,Hz are the x, y, z components of H efr, and can be derived by:

dF

d M y

Hy = Hs\n(0 - qi) (2.7)

d F

d M ,+ Heos(0 - <p)

Therefore, the torque equations become:

1 drriy

y d t= — [H cos(0 — <p) + {Ha — Hd) * cos2B]my

^ = [H cos(0 -<p) + (Ha - Hd) * cos(29)]mx (2 .8)

For the uniform precession, the my and mz can be written as:

my = m yQe l2TCft+<py

m z = m z0e i 2 n f t+ < p z (2.9)

The precession frequency f of spin wave can be derived by substituting Eq.

(2.9) into Eq. (2.8):

2tt/ = yV[W cos(0 - <p) + (Ha - //d) * cos(28)][H cos(0 - <p) + {Ha - Wd) * cos20]

(2 .10)

23

In terms of Gilbert damping parameter a, one has to incorporate the

damping term in Eq. (2.6), and the uniform magnetization precession can be

described by an oscillating term times the decay term where

a y [ ( H cos(0 - <p) + ( tfa - tfd) * cos20 ) + ( / / cos(0 - <p) + ( t fa - tfd) * cos(20))]

T “ 2(1 + a 2)

(2 .11)

is the lifetime of magnetization precession. Therefore, given the precession

frequency f and lifetime r as a function of magnetic field H, we can derive

some physical parameters of the magnetic material, such as magnetic

anisotropy and effective Gilbert damping parameter aeff.

2.2.3 Experimental set-up

Figure 2.7 displays the schematic TRMOKE setup used in this

dissertation. TRMOKE measurements are performed in a pump-probe setup

using pulsed Ti.sapphire laser (RegA-9000 amplifier laser system) with a

pulse duration of 200 fs and a repetition rate of 250 kHz. The wavelength of

pump (probe) pulses is 400 nm (800 nm). A modulated pump pulse beam

with a fluence of 0.16 mJ/cm2 is focused to a spot ~ 1 mm in diameter on the

sample to excite the magnetization precession, and the transient Kerr signal

is detected by a probe pulse beam which is time-delayed with respect to the

pump. The focus area of probe beam is with a diameter of 0.7 mm, which is

24

smaller than that of pump beam; so that the intensity ratio of the pump to

probe pulses is set to be about 6:1. A variable magnetic field H is applied

using a superconducting magnet produced by Oxford Instruments. The probe

signal is send to a Lock-in amplifier produced by Stanford Instruments.

Polarizer (P)

Analyzer (S)

Electromagnet

Figure 2.7 Schematic of TRM OKE experimental set-up.

25

2.3 Magnetization-induced second harmonic generation

MSHG is a nonlinear optical technique that probes the interface or

surface magnetization [26], and sometimes it is called nonlinear MOKE. The

MSHG signal can be generated in magnetic materials in which both space-

inversion and time-reversal symmetry are simultaneously broken. The broken

time-reversal symmetry results from the presence of magnetization, while the

broken space-inversion symmetry occurs only at the surface or interface of

centrosymmetric materials. Figure 2.8 depicts the different probing regions of

MOKE and MSHG in La0.6 7 Sr0.3 3 MnO3 magnetic thin film capped with

PbZr0 5 2Tb 4 8O 3 non-magnetic layer. MOKE probes the bulk magnetization,

while MSHG probes the interface magnetization.

MSHG MOKE

Interface

LSMO

Figure 2.8 Schematic of MSHG and MOKE measurements on PZT/LSM O thin film

26

2.3.1 Second harmonic generation (SHG)

SHG signal arises from the nonlinear polarization P(2oS) induced by an

incident laser field E(co), where co is the angular frequency of light. This

polarization can be written as an expansion in E(co):

P(2o>) = x {2)E{ai)E(a>) + x (Q}E(o))VE(a)) + - (2.12)

The lowest-order term in Eq. (2.12) describes an electric dipole source, with

* (2) as a third-rank susceptibility tensor. Symmetry considerations show that

this contribution is zero in a centro-symmetric medium, thus limiting electric

dipole radiation to the surface or interface where space-inversion symmetry is

broken. The second term in Eq. (2.12) describes the bulk SHG in terms of the

much smaller electric quadrupole-like contributions, where * (<2) is fourth-rank

susceptibility tensor. However, because of the large volume difference

between interface and bulk this does not necessarily mean that the total bulk

second-harmonic signal is negligible. Moreover, electric field induced SHG is

also generated in the bulk of centrosymmetric materials, and this effect can

be significant since the quasi-static electric field can reach very high values.

[27]

2.3.2 Nonlinear magnetic susceptibility

The presence of magnetization will not break the space-inversion

symmetry of the bulk, but can lower the surface and interface symmetry

modifying the nonlinear susceptibility. As a result, the surface or interface

27

second harmonic polarization can be simplified to one third rank tensor with

different components:

Pt( 2a>) = <JCijkw ( ± M ) + Xijk(-H±M))Ej(a>)Ek(a>) (2.13)

where i, j, k=x, y or z indicate the directions of longitudinal, transverse and

polar components, * ijk(+) and represent the tensor components even

and odd in M, respectively. The odd component ^jjk(_) changes sign upon

magnetization reversal and therefore give rise to the magnetic asymmetry in

the MSHG response.

The symmetry of the nonlinear susceptibility is determined by the

symmetry of the particular surface under consideration. Some components of

X are non-zero, if they satisfy the same space symmetry. Therefore, the non­

zero elements of x can be obtained from invariance of x under symmetry

operations:

Xi \k = Tii'VTkk'^i'j'k' (2-14)

where T is the transformation matrix for each symmetry operation, and

summation over repeated indices is implied. Take the magnetized (001)

surface with cubic symmetry for example: With M aligned along [100] (x)

direction, the two independent symmetry operations are reflection about the

x-z mirror plane, (x ->x, y ->-y, and Af-> -M) and reflection about y-z mirror

plane, (x ->-x, y ->y, and M). By utilizing Eq. (2.14), one can obtain the

non-vanishing elements. Xxxz> Xxzx> Xyyz> Xyzy> Xzxx> X z y y 3nd Xzzz< which are

28

even terms, x Xxy> Xxyx> Xyxx> Xyyy> Xyzz> X z y z X z zy odd terms. Table

2.2 summarizes the non-zero terms in different cases.

Table 2.2 Nonzero tensor elements of the nonlinear susceptibility tensor * (2) for the

(001) cubic surface, the magnetization parallel to the x-axis (longitudinal

configuration), y-axis (transversal configuration) and parallel to the z-axis (polar

configuration). and denote the even and odd elements with magnetization.

Other non-zero elements with mixed polarization of s and p are not included. For

simplicity of notation the tensor components are indicated by their indices only.

InputPolarization

OutputPolarization

Longitudinal

J lf//x

yzz, yxx P s

zzz, xzx xxz, zxx

P P

zyy s Pyyy s s

Transverse

H f//y

P sxxz, xzx zxx, zzz

zxz, zzx XXX, xzz

P p

zyy xyy s ps s

Polar

M / / z

yxz, yzx p sxxz, xzx zxx, zzz

p p

zyy s ps s

29

2.3.3 Experimental consideration

The MSHG response depends not only on the geometry of

magnetization in relation to the incident light and sample plane, but also on

the polarization state of incident and outgoing laser beam. Table 2.2

categorizes the non-zero susceptibility tensor components obtained for

different polarization configurations. For the MSHG measurements on

PZT/LSMO heterostructure, the s-in s-out polarization configuration is used.

The odd susceptibility tensor element is xyyy(Mx) as a function of

magnetization along longitudinal direction, the even term comes from xzyy as

we rotate the detection analyzer 5 degree away from perfect s-out to

introduce a non-magnetic background for improving the magnetic contrast

(see below). The second harmonic signal comes from electrical polarization of

PZT mainly contributing to the term Xzzz and is eliminated in this optical

polarization configuration.

The measured intensity in a fixed experimental geometry with opposite

magnetization direction can, in general, be written as a sum of effective

tensor components:

IshH 2 co) oc \XeffW ± Xeff(~)\2 (2.15)

where T e //(+) and * e / / (_) are the linear combinations of even and odd tensor

components with Fresnel coefficients a iJk\

X e f f ^ l i , j . k & i j k X i j k (2-16)

30

The magnetic contrast A for MSHG measurement is defined as

where l(±M) is the signal intensity received by the detector when the

macroscopic magnetization M is aligned with external magnetic field in either

positive or negative magnetic field direction (Fig. 2.9). This contrast is

normalized to the total SHG intensity, and does not depend on the intensity of

the fundamental light. Moreover, the magnetic contrast A obtained by MSHG

is linear with M in the probed volume, if magnetic component is much smaller

than non-magnetic component in signal intensity as in our case [28].

0.36(+M)

- - I (-M)

0.28

-1000 0 1000Magnetic Field (Oe)

Figure 2.9 MSHG hysteresis loop from PZT/LSM O thin films indicating l(-M) and

l(+M) used to determine the magnetic contrast A from Eq. (2.17).

31

2.3.4 Experimental set-up

MSHG experiments are performed with a Ti:sapphire laser (RegA-

9000 amplifier system) generating 200 femtosecond pulses with 4 pJ energy

at 250 kHz repetition rate and 800 nm wavelength. The attenuated laser

beam (50 mW) with s-polarization is focused to a 200-pm diameter spot on

the sample at an angle of incidence of 40°. A small S-polarized MSHG signal

(400nm) is generated in the direction of the reflected laser beam, and is

detected with a high signal-to-noise ratio photomultiplier produced by

Hamamatsu company. Filtering with prism is required to separate the MSHG

light from the fundamental laser beam. The detection scheme A and B are for

MSHG and MOKE measurements respectively.

Red-psfiltei Prism

Lens Lens Photodiode

AnalyzesPolarizer

Figure 2.10 Schematic set-up of MSHG (A) and MOKE (B) measurement.

32

Chapter 3

Time resolved studies of spin dynamics in FePt thin flims

The coherent spin precession in L1o ordered FePt thin films is

investigated with TRMOKE technique. The perpendicular magnetic anisotropy

Ku and intrinsic Gilbert damping oq of the material systems are derived from

spin dynamics studies. Furthermore, the quadratic dependence of Ku and ao

on spin-orbit coupling strength £ is rigorously observed for the first time,

where £ is continuously controlled through chemical substitution of Pt with Pd

element. These results are consistent with first-principle calculations, which

also show that other leading physical parameters remain unchanged through

doping.

Moreover, a linear correlation between do and electron scattering rate

1/re is experimentally observed through modulating the anti-site disorder c in

the L10 ordered FePt thin films. The variation of c mainly affects re, while

other leading parameters remain unchanged. Moreover, as c decreases, the

perpendicular magnetic anisotropy increases and the Landau g factor exhibits

a negative shift due to an increase in orbital momentum anisotropy.

33

3.1 Introduction

Ultrafast magnetization precessional switching in ferromagnets utilizing

magnetic field pulses, spin polarized currents, and ultrafast laser pulses [29-

33] is currently a popular topic due to its importance in magnetic information

storage and spintronic applications. The uniform magnetization precession

can be well modeled with LLG equation, where the Gilbert damping

parameter a determines the spin relaxation time [34] and is crucial for device

performance [9-12]. The extrinsic Gilbert damping is due to nonlocal spin

relaxation, such as spin pumping and two-magnon scattering, which can be

tuned by artificial substrates, specially designed buffer and coverage layers

[35-41],

The intrinsic Gilbert damping a0 describes the energy flow rate from

spin to electronic orbital and phonon degrees of freedom through electron

scattering, and has been studied in various theoretical models [42-50]. The

breathing Fermi surface model [42] and torque-correlation model [43] based

on first-principle band-structure calculations qualitatively match or0 in soft

magnetic alloys, such as Fe,Co and Ni [51-56]. Moreover, contributions to do

can be categorized based on intraband and interband transitions [43, 46]. The

damping rate from intraband transitions scales linearly with the electron

relaxation time re and exhibits conductivity-like behavior. In contrast, the

damping rate from the interband transition is proportional to the electron

scattering rate 1/re, and consequently exhibits resistivity-like behavior.

Moreover, contributions of intraband (interband) transitions to damping are

34

predicted to play a dominant role in low (high) temperature regions, and be

proportional to £3 ( f 2).

Perpendicular magnetic anisotropy (PMA) defines the low-energy

orientation of the magnetization M (easy axis) normal to the film plane, as well

as the stability of M with respect to external fields, electric currents, and

temperature-induced fluctuations. Materials with large PMA are good

candidates for high-density memory and spintronics devices [8-10], as they

exhibit promising thermal stability and allow going beyond the

superparamagnetic effect, giving access to magnetic media with smaller

magnetic domains [7,8]. Moreover, PMA-based magnetic tunnel junctions

provide high tunneling magneto-resistance ratio [9-11], and low critical current

for spin-transfer-torque switching [12]. Considering these advantages,

tailoring the PMA of magnetic materials as well as elucidating its physical

origin becomes important for further technological advancements.

PMA can be related to the spin-orbit-coupling (SOC)-induced splitting

and shifting of electronic states that depend on the magnetization direction,

which results from the simultaneous occurrence of the spin polarization and

SOC [57-59]. An effective method to achieve significant PMA utilizes

appropriate combinations of 3d and heavier Ad, 5d elements, which merge

the large magnetic moment and magnetic stability of 3d TMs with the strong

SOC of Ad, 5d TMs [60-64], L10 ordered FePt material is one prominent

example [65-69]. In addition, Pt acquires a sizable spin polarization in contact

with Fe and contributes significantly to the PMA, due to its large SOC.

35

3.2 Samples and characterization

Figure 3.1 shows the schematic of atomic structure for the two series

of FePt thin films where the L10 ordered structure consists of alternate

stacking of Fe and Pt atomic planes along the face-center-tetragonal (fct)

[001] direction. In the first series (left, SOC dependent), continuous

substitution of Pt by Pd leads to a gradual decrease of £ in L1o ordered

FePd(i-X)Ptx alloy, while other leading physical parameters remain almost

unchanged. That is because the element Pt falls into the same group as Pd in

the periodic table, and £(Pt) is stronger than £(Pd). In the second series (right,

chemical order dependent), the anti-site occupation of atoms is controlled

with proper growth temperature, which mainly affects the electron scattering

rate 1/reand the hybridization effect between Fe and Pt atoms.

The two series of L10 FePt(Pd) alloy films are deposited on single

crystal MgO (001) substrates by magnetron sputtering. The FePt(Pd)

composite target is fabricated by placing small Pt (Pd) pieces on a Fe target.

The base pressure of the deposition system is 1.0*1 O'5 Pa and the Ar

pressure is 0.35 Pa. During deposition, the substrates are kept at 7"g=620 °C

for the SOC dependent study and at different temperatures Tg for the

chemical order dependent study, and the rate of deposition is about 0.1 nm/s.

After deposition, the samples are annealed in situ at the same temperature as

for the growth Tg for 2 hours. The film thickness is determined by x-ray

reflectivity to be 17 ± 1 nm (22 ± 1 nm).

36

t • • • • • • •• • • • • • •• • • • • • » Fe

[001] • • • • • • • ^• • • • • • •Perfect L1(0) ordered Fe/Pt alloy

/

• • • • • • •• • • • • • § Fe• • • • • • • • • • • • •

• • • • • • •• • • • • • « Fe• • • • • • •• • • • • • • Pt

• • • • • • • • Pd • • • • • • •FePt(x)Pd(l-x)

modulating spin-orbital coupling Varied chemical order

Figure 3.1 Schematics of atomic structures for the two series of FePt thin films.

Structure characterization of FePt(Pd) samples is performed with X-ray

diffraction measurement with Cu Ka radiation, as presented in Fig. 3.2. Only

fct (001) and (002) peaks of FePt(Pd) are observed in the spectrum along

with other peaks from MgO substrate, which indicates the L1o ordering in the

FePt(Pd) alloys. The peak positions do not shift with different Tg or Pd doping

level (1-x), which indicates that the lattice constant varies by less than 1.0

percent and tetragonal distortion of lattice is not affected by either doping or

modulation of Tg. The anti-site disorder percentage is derived from

°Q2' meas\ /omeas (3.1)('001)' 002' mi

37

where S is the degree of chemical order, /001 and /002 are integrated

intensities of fct (001) and (002) peaks, and (1 - c ) is the probability of the

correct site occupation for either Fe or Pt atoms in such L l0 ordered alloy

system [70]. U o o i / I oo2 )c a ic *s calculated to be 2.0 for the perfect ordered film

with thickness ranging from 11 to 49 nm [70], Thus for the Pd doped FePt thin

films, S is about 0.8, indicating a good suppression of disordering defects in

these films and about 90% of atoms stay at the correct sites. Moreover, c as a

function of Tgfor both 17-nm-thick and 22-nm-thick films is shown in Fig. 3.3.

Higher Tg leads to monotonous decrease of the anti-site disorder c in FePt

alloys, as depicted by the insets in Fig. 3.3.

38

X—1

x=0.5x=0.25

&toc<D-4—*

Qa : x fct (002)fct (001)

6040 5020 3020 (deg)

S

(/)c0coa:X

(b) fct (001)

Tg=720 °C f jfct (002)

u| Tg=680 °C P A

Tg=620 °C I I Aa Tg=580 °C

MgO J L1 1 i ......

20 4020

60

Figure 3.2 Structure characterization results of FePtxPd{1.x) thin films (a) and FePt

thin films with different Tg (b) by X-ray Diffraction measurements.

39

o

1 6

12

8

• • • • • • •• • • • • • * FeJ

• • • • • • •• • • • • • • Pti• • • • • • •

2 2 nm F e P t 17 nm F e P t

• • • • • • • •! • • • • • • f t •

- ' • • • • • • •Pt J• • • • • • • •

• • • • • • •

480 540 600 660Tg(°C)

720

Figure 3.3 Anti-site disorder percentage c in L10 ordered FePt thin films as a

function of growth temperature Tg.

40

"(a) X=1

“ (b)x=0.5

"(C)

— 0 . =ox=0- - 0 = 90

H (T)0 0006

0.0004

0.0002

0.0000 r

- 0.0002

-0.0004

-0 0006- 3 - 2 - 1 0 1 2 3

Magnetic field H (T)

Figure 3.4 Static magnetic hysteresis loops measured by vibrating sample

magnetometer (VSM) (a)-(c) for FePtxPd(i.X) thin films. The black (red) curves

represent out-of-plane (in-plane) results. VSM results for FePt thin films with

different Tg (d).

41

Figures 3.4(a)-(c) display the out-of-plane and in-plane magnetization

hysteresis loops of FePtxPd(i-X) thin films for x = 1.0, 0.5, and 0, respectively.

As shown in Fig. 3.4(a), for x =1 (L l0 FePt) the out-of-plane hysteresis loop is

almost square-shaped with coercivity Hc = 0.44 T; but it is hard to reach the

saturated magnetization with in-plane magnetic field, indicating the

establishment of high PMA. With decreasing x, Hc decreases due to the

reduced PMA. For x = 0 in Fig. 3.4(c), Hc approaches zero and the out-of­

plane and in-plane hysteresis loops almost overlap with each other,

indicating weak PMA. The PMA therefore increases with increasing

concentration of Pt in FePd(i.X)Ptx alloy thin films. From the experiments, the

saturation magnetization Ms for all samples is 1100 emu/cm3 within 10%

relative error, close to the bulk value of L1o FePt [71].

Figure 3.4(d) displays the out-of-plane magnetization hysteresis loops

for 22-nm-thick films with Tg = 580 °C, 620 °C, and 680 °C, and the in-plane

hysteresis loop for the sample with Tg = 620 °C. The out-of-plane hysteresis

loops are almost square-shaped with coercivity Hc = 0.3 T; but it is difficult to

reach the saturated magnetization with in-plane magnetic field, indicating the

establishment of high perpendicular magnetic anisotropy Ku. From the

experiments, the saturation magnetization Ms for all samples is determined to

be 1100 emu/cm3 and remains unchanged as a function of growth

temperature and disorder c. Moreover, the magnetization is not fully saturated

with H = 2T applied along the easy axis, indicating the existence of multiple

magnetic domains at lower magnetic fields.

42

3.3 Magnetic anisotropy and damping in FePtxPd (1.x) thin films

Figure 3.5 shows TRMOKE results of FePd(i.X)Ptx films with various x

at H = 5 T. A uniform magnetization precession is excited as manifested by

the oscillatory Kerr signal 0k, while the magnetic damping is indicated by the

decay of precession amplitude as the time delay increases. As shown in Fig.

3.5, the measured Kerr signal can be well fitted by the following equation

0k = a+b*exp{-tAo)+A*exp(-tlT)s\r\(2jxft+q>) (3.2)

where parameters A, r, f and <p are the amplitude, magnetic relaxation time,

frequency, and initial phase of the magnetization precession, respectively.

Here, a, b, and to are related to the background signal owing to the slow

recovery process after fast demagnetization by laser pulse heating. It is well-

demonstrated in Fig. 3.5 that the spin precession frequency and magnetic

damping effect become larger and stronger for higher doping level x with the

same H.

In order to derive the PMA and ao for FePdi.xPtx samples at different

doping levels, magnetic field-dependent TRMOKE measurements are

performed. As shown by the TRMOKE results of x = 0.82 and x = 0.5 samples

in Fig. 3.6(a) and 3.6(b), spin precession frequency and relaxation time

increases as H increases. The field dependence of frequency can be

understood by the enhancement of the effective field due to larger H. As for

the damping, the strong magnetic field suppresses the extrinsic part of Gilbert

damping, such as caused by dephasing, and a is governed by intrinsic

43

contribution at higher magnetic fields. The H dependences of the precession

frequency f and relaxation time r for some typical samples are displayed in

Figs. 3.7(a)-3.7(b). W e note that f can be widely tuned from 15 GHz to 340

GHz by varying the magnetic field and doping level, and r decreases by two

orders of magnitude when Pd atoms are all replaced by Pt ones and reaches

about 7.0 ps for x=1 (L1o FePt).

From the LLG equation with cr«1.0, the dispersion equation is

2TTf = ><H,H2)1/2, (3.3)

where H i= H c o s ( 0 h - 0 ) + / - /kc o s 20 and /-/2= /- /c o s (0 w-0 )+ H k C O s 2 0 , Hk=2KJMs-

4ttMs with uniaxial magnetic anisotropy constant Ku, gyromagnetic ratio y

and 9h = 45°. The equilibrium angular position G of the magnetization satisfies

the following equation sin20 = (2/-///-/K)sin(0H-0). The measured field

dependence of f can be well fitted by Eq. (3.3), as shown in Fig. 3.7(a). With

the measured Ms of 1100 emu/cm3, and the g factor fitted to be 2.1 ± 0.05 for

samples with different x, Ku is derived from the dispersion and displayed in

Fig. 3.8(a) as a function of doping level x. For a « 1 .0 , aro can be in principle

extracted by fitting the measured field dependence of t with t =2/aroY/(Hi + H2)

and using the fitted values of g and HK, as shown in Fig. 3.7(b). The deviation

between fitting and data points at lower magnetic field is due to the

appearance of extrinsic Gilbert damping, which mainly comes from the

dephasing dynamics caused by inhomogeneity of PMA in thin films, as the

magnon scattering is less effective for perpendicularly magnetized samples

44

[72], The measured «o exhibits a sensitive dependence on the chemical

substitution in Fig. 3.8(b).

X=1

x=0.82

x=0.5zj

x=0.25CD

=0.15

30 4010 20Delay t (ps)

Figure 3.5 TRM OKE data with different doping level of FePtxPd(l-x) under magnetic

field H= 5T . The solid lines are fitted curves.

45

x= 0.82 H= 5 T

H= 4 T

H= 3 TCD

H= 2 T

30 4010 20Delay t (ps)

zsCD

o'

20 40 60 80Delay t (ps)

Figure 3.6 TRM OKE results of FePt0.82Pd0.18 (a) and FePt0.5Pd0.5 (b) under different magnetic fields H. The solid lines are fitted curves.

* X= 0 5 H= 5 Tj / V \ A a a a a a a / w w

i/ V V v v ^ ^

H= 2 T

46

0.3

0.2N

XI -

X=1x=0.9x=0.75x=0.5x=00.0

0 2 4 6H (T)

0.6• x=0.9 ▲ x=0.75 ▼ x=0.5♦ x=0

0.3

0.04

0.000 2 4 6

H(T)

Figure 3.7 The dependence of precession frequency f (a) and lifetime x (b) on

external magnetic field H fo r typical samples. The solid lines refer to fitted results.

47

^ 4COE ,0^5)0

r -O 2T—zs

00.0 0.5 1.0

X

0.10

0.05

0.00

Figure 3.8 PMA (a) and a0 (b) as a function of doping level x of FePtxPd(1.x).

Mb)

1.00.50.0

n

- /— i

•■ ■ ■

I i . i

48

0.6

>d) 0.4

0.2

15>

10

5 5

CD3.0ZL

U) 1.5F

0.01.5

<0 1.0O

0.5

-(b)

(a)

- (c)

'(d)

0.0

■I

__L _

0.5x

FePt/Pd

■ ■

1.0

Figure 3.9 Spin-orbit coupling strength f (a), bandwidth tV (b), spin moment mson Fe

or Pt/Pd site (c) and measured c/a ratio (d) as a function of doping level x of

FePtxPd(i.x).

49

To investigate the spin-orbit interaction in FePdi-xPtx doping system, ab

initio density functional calculations are performed to quantitatively determine

the key physical parameters [73, 74], The calculated f , bandwidth W and

magnetic moment ms at Fe or Pd/Pt site as a function of doping level x is

derived and depicted in Figs. 3.9(a)-3.9(c). Figure 3.9(a) shows that f at

Pd/Pt site changes from 0.19 to 0.57 eV when x varies from 0 to 1.0. For Pt,

Pd, and Fe atoms, £ is 0.6, 0.20, and 0.06 eV [73, 74], respectively, and

therefore the effect of Fe atoms is negligible compared with those of Pd and

Pt atoms. The change of spin moment and bandwidth are negligible when x is

tuned from 0 to 1, as shown in Fig. 3.9(b) and 3.9(c). The Pt/Pd atom

acquires spin polarization around 0.27 (pe/atom) due to the proximity effect.

Moreover, the lattice constant remains unchanged for different x from both

XRD results and theoretical structure analysis of L l0 FePdPt alloys, and c/a

ratio is around 0.94 as shown in Fig. 3(d). Furthermore, the density of states

at Fermi level D(EF) is found to change from 2.55 to 2.39 per atom per eV for

x varying from 0 to 1.0. Standard electron-phonon scattering calculations are

performed with the phonon Debye model and static limit of Lindhard’s

dielectric function. The electron-phonon scattering rate 1/re varies from 1.34

to 1.33 ps"1 for x changing from 0 to 1.0. Since other leading parameters

remain unchanged, the variation of PMA and ar0are mainly controlled by SOC

strength through different doping level x.

Figure 3.10 (a) displays the measured values of ar0as a function of£,

which agrees with the calculated values from first principles. The results

50

rigorously prove the theoretical prediction of the £2 scaling of ato for the first

time, indicating that ar0 at high temperatures is mainly caused by interband

contribution. Moreover, the electronic-scattering-based model of FM

relaxation is therefore proven to be applicable for ao in L10 FePdPt ternary

alloys.

Figure 3.10(b) shows the £ dependence of PMA derived from

TRMOKE measurements (black squares). The magnetic anisotropy arises

from second-order energy correction of SOC in the perturbation treatment.

Therefore, the PMA is roughly proportional to both f and orbital magnetic

moment when f is smaller than the exchange splitting. Since the orbital

moment is also proportional to ( within perturbation theory; the enhanced

PMA at higher x is attributed to a larger £ of Pt atoms compared with that of

Pd atoms and a quadratic scaling law of PMA with f is expected, as observed

from the experimental results. Moreover, first-principle calculations are

performed [74] and the calculated values of PMA are larger than the

measured ones at higher x (red spots in Fig. 3.10(b)). Several reasons exist

for the discrepancy, such as slight differences in the lattice structure between

experiment and theory, and thermal fluctuations and chemical disorder are

not accounted for in the theoretical calculations. Hence, there is potential for

further enhancing PMA by SOC tuning.

One approach to enhance the SOC dependence of PMA is increasing

the chemical order of FePdi.xPtx alloys, as the large PMA values in the

theoretical study are based on an ideally ordered lattice. To estimate the

51

influence of disorder, we perform TRMOKE experiments on the samples with

higher chemical ordering S = 0.85, where the correct site occupation of atoms

is increased to 92% from rFe = r Pt Pd) = 90% (S = 0.8). The measured PMA

values, indicated by the blue rhombus symbols in Fig. 3.10(b), show a slight

increase for x=0.5 (£ = 0.26) and a noticeable enhancement for x=1 (£ = 0.57).

This indicates that the chemical ordering of FePdi-xPtx alloys has a strong

effect on PMA by spin-orbit interaction, as the disordering can affect the

electronic band structure [75]. Also at lower temperature (20K) the PMA

values of FePdi.xPtx alloys can be more enhanced by SOC strength due to the

suppression of thermal spin fluctuation, as shown in Table 3.1. Furthermore,

other approaches such as modulation of tetragonal distortion and chemical

substitution of 3d transition metal elements in alloys may also be utilized for

SOC tuning of PMA.

52

E.o13>Q)

O

13

10■ measured • calculated8

6♦ chem ordered

4

2

0

0.0 0.3 0.6SOC (ev)

measured0.10

s° 0.05

0.00

0.0 0.3 0.6SOC $ (ev)

Figure 3.10 Measured and calculated cro (a) and PMA (b) as a function of SOC

strength. The solid line is a quadratic fit obtained from first principle calculations.

53

FePto.5Pdo.5

Temperature (K) 9 Hk(T) Ku (107erg/cm3)

20 2.13 2.17 2.07

200 2.10 2.08 1.914

FePto.7 5 Pdo.25

Temperature (K) 9 « k (T) Ku (107erg/cm3)

20 2.20 3.96 3.097

200 2.18 3.80 2.849

FePt

Temperature (K) 9 Hk (T) Ku (107erg/cm3)

20 1.97 7.30 5.012

200 2.00 6.96 4.587

Table 3.1 Landau g factors and Ku derived from TRMOKE measurements for

different doping level x at 20K and 200K.

54

3.4 Magnetic anisotropy and damping in FePt thin films with

different chemical order

Figure 3.11(a) shows the TRMOKE results of FePt thin films with

various c at H = 6.5 T. The uniform magnetization precession is demonstrated

by the oscillatory Kerr signals 0K, while the magnetic damping is indicated by

the decaying precession amplitude as the time delay increases. The

measured 6k can be well fitted by Eq. (3.2). It is well demonstrated in Fig.

3.11(a) that the spin precession frequency is larger and magnetic damping

effect becomes weaker for lower c with the same H.

In order to obtain the PMA and or for FePt samples with different c,

magnetic field (H) dependent TRMOKE measurements are performed. Figure

3.11(b) shows the measured results of the 22-nm-thick sample with Tg = 680

°C. It can be seen clearly that the precession period and relaxation time vary

as H increases. The fitted fas a function of H for different c are plotted in Fig.

3.12(a). We note that f can be tuned from 240 GHz to 335 GHz by varying H

and x. The measured H dependence of f can be well fitted by Eq. (3.3), as

shown in Fig. 3.12(a), and we thus obtain Ku and g. Ku decreases from 6.0 to

3.2 (107 erg/cm3) and the Landau g factor also displays a shift from 1.87 to

2.24 as the anti-site disorder percentage c increases, as shown in Figs.

3.13(a)-3.13(b).

55

13TO,*<D

(a)c= 4.0%

10

c= 8.6%

c= 10.5%

c= 13.8%

20 30t(ps)

c= 4.0% H= 6 T

H= 5 T=3d

H = 4 T

H= 3 T

20 30t(ps)

Figure 3.11 TRMOKE data from 22-nm thick FePt films, grown at different Tg with

applied magnetic field H -5T (a), and with Tg =720 °C measured at different H (b).

The solid lines refer to fitted results.

56

0.35c= 4.0% c= 8.6% c— 10.5% c= 13.8%0.30

NII -- 0.25

0.202 64

H (T)

▼ c= 13.8% a c= 10.5% • c= 8.6%■ c= 4.0%

0.20

0.15

4)

0.10

0.05

2 4 6H (T)

Figure 3.12 The dependence of spin precession frequency f (a) and effective Gilbert

damping aeff (b) on H for FePt thin films with different c. The solid lines are fitted

results.

57

EoE>a>

h-o

=3

(a)

■ 22nm FePt - 17nm FePt

0 6 12 Anti-site percentage c

18

2.2

oo£ 2.0o>

6 12 180Anti-site percentage c

Figure 3.13 PMA Ku (a) and Landau g factor (b) as a function of anti-site defect

concentration c for FePt films with thicknesses of 17 nm and 22 nm.

58

Figure 3.13(a) shows that the Ku maintains high values from 3.2 to 6.0

(107 erg/cm3), and gradually increases with smaller c, which is consistent with

other experiments, as well as the theories [76-79]. The Ku in FePt alloys

results from the simultaneous occurrence of the spin polarization and large

SOC [74], The smaller c indicates that more Pt atoms become nearest

neighbors of Fe, which results in stronger hybridization between Fe and Pt

atoms. Consequently, Pt acquires larger spin polarization and orbital moment

due to the Fe-Pt hybridization in the sample with higher chemical order, and

contribute significantly to the Ku. As a result, the orbital momentum anisotropy

[78, 79] and hence Ku increase with decreasing c.

Figure 3.13(b) reveals a gradual decrease of Landau g factor with

smaller c. The g factor describes the proportionality of angular momentum

and magnetic moment for the individual spins, which also affects the dynamic

response of a magnetic film. In itinerant electron systems, the g factor may be

written as

(3.4)y e <S'>+</-> v '

where ps (ul) denotes the magnetic moment from the spin (orbit)

component. For a symmetric crystal lattice, the orbital motion of d electron is

quenched by crystal field effect, i.e., <L> = 0. Nevertheless, the orbital

contribution to the magnetic moment is nonzero, thus the g factor is equal to

2*(1+ piJps) and is typically greater than two in itinerant electron system.

However, as c decreases, the enhanced hybridization between Fe and Pt

59

restores the orbital momentum due to the large SOC strength of Pt [78], and

raises the orbital momentum anisotropy [78, 79]. This would lead to <L> * 0,

and s f 2 * ( 1 - g J fJ s ) , indicating a negative shift of g factor relative to the value

of two. Such negative shift of g factor is also observed at surface or interface,

where the orbital momentum is not entirely quenched due to the broken

symmetry [80].

Using the fitted values of r, we determine a with

a , „ = a„ + a „ = (3.4)

where arex is the extrinsic contribution to Gilbert damping. As shown in Fig.

3.12(b), the value of aeff gradually decreases with higher H and saturates at

higher fields. The decreasing trend of creff with H is attributed to the

suppression of dephasing dynamics among magnetic domains, since multiple

domains exist at low fields as indicated in Fig. 3.12 (b) and the magnon-

magnon scattering is less effective for perpendicularly magnetized samples

[72]. The saturation values of areff at higher fields are therefore used to

approximate the intrinsic Gilbert damping ao [73]. Figure 3.14(a) shows cto as

a function of c at 7=200K.

60

0.21

o>■£= 0.14Q .E(0■O■Ca>

S 007

(a)22nm FePt 17nm FePt

0 6 12 18 Anti-site percentage c

0.09

0.06o>

0.03

22nm FePt 17nm FePt

0 6 12 18 Anti-site percentage c

Figure 3.14 a0 (a), and or</g (b) as a function of anti-site defect concentration c for

FePt films with thicknesses of 17 nm and 22 nm. The solid line refers to least square

linear fit.

61

The key finding here is that or0 gradually increases with larger c. Since

the scattering rate 1/re is enhanced with more impurity scattering sites c

according to scattering theory [48-50], the positive correlation between aro and

1/je qualitatively matches the resistivity-like behavior where or0 is governed by

interband transitions [43, 46]. In the spin-orbit coupling torque correlation

model for interband transitions [43, 46],

a„ « g ( f /M 0 2A« (3.5)

where W is the d-band width, and f the spin-orbit coupling strength. It is

demonstrated that anti-site disorder in L l0 ordered alloys smoothens the

density of states [81]. The calculated D(Ep) increases less than 5% as c

increases, and W remains almost the same [75], when 0<c<15% as in our

case. In our previous work [82], we investigated the anomalous Hall

conductivity of L1o ordered FePt films with different ordering. We found that

the resistivity independent term (bo) of anomalous Hall conductivity remains

almost unchanged with x when x is small, indicating that the variance of spin-

orbit coupling strength £ in our samples is negligible. The g factor increases

by 1 9 % with larger c as shown in Fig. 3(a), due to the modulation of orbital

momentum anisotropy. The lattice distortion and Ms remain almost

unchanged as demonstrated by structure characterization and VSM

measurements. A similar trend of aro with c is observed for samples with two

different thicknesses, which indicates that surface and strain effects on

damping can be ruled out. Therefore, we attribute the significant increase of

62

aQ by more than a factor of three, as revealed in Fig. 3.14(a), to the

enhancement of 1/re as a result of the increasing c. To separate the effect of

g on the damping, we further plot aotg as a function of c in Fig. 3.14(b). We

observe approximately a linear correlation between the two variables. The

exchange of different elements in L10 ordered alloy film leads to scattering of

itinerant electrons through local spin dependent exchange potentials. Since

the cross-section and strength of individual scattering events remain

unchanged in weak scattering regime [48], the linear correlation indicates that

do is proportional to 1/re, where the damping process is dominated by

interband contribution. The damping process can be considered roughly as

the decay of a uniform mode magnon into an electron spin-flip excitation.

Therefore, the anti-site disorder works as spin-flip scattering center for

itinerant electrons transferring spin angular momentum to the lattice via SOC.

Moreover, complete suppression of the anti-site defects might lead to a

remnant cro, where electron is mainly scattered by phonon instead of impurity.

In order to check for the intraband contribution to cro, which may become

significant with low 1/re, TRMOKE measurements are carried out for the 17-

nm-thick FePt film with the lowest c=3% at low temperature. Figure 3.15(a)

shows that the frequency and decay rate of coherent spin precession at H=5T

varies slightly with temperature (from 20 K to 200 K). The field dependent

frequency and decay rate at 20 K, as shown in Fig. 3.15(b), are analyzed to

obtain a0. It turns out that a0 remains almost unchanged (from 0.053 ± 0.013

to 0.054 ± 0.013) when temperature decreases from 200 K to 20 K, despite a

63

significant change in the electron-phonon scattering rate [43, 46]. The

temperature independent behavior of cro is due to the fact that electron-

impurity scattering dominates the scattering events. In Chapter 3.3, we

calculated the electron-phonon scattering rate from first principles to be

approximately 1.33 ps'1 for FePt at 200 K. The electron-impurity scattering

rate must be considerably higher, consistent with the values of electron

scattering rate from Ref. [46] where the interband contribution to cro dominates.

Moreover, the weak temperature dependence of cro indicates that 3% anti-site

disorder is still too high to observe the conductivity-like behavior for intraband

contribution to cr0, which is consistent with theoretical prediction [49]. Further

investigations at low temperature with fewer impurities are necessary to

elucidate the relationship between cr0 and 1/re governed by intraband

transitions.

64

0.3 (a)

N 0.2

0.1 - •

0

Frequency Decay rate

100 T (K)

0.3

0.2 £

0.1

200

0.30.3

■ Frequency • Decay rate0.2

NXI—

3 64 5H (T)

Figure 3.15 The dependence of frequency f and decay rate 1/ron temperature T

with applied field H =5T (a), and as a function of H at 20 K (b). The solid lines refer to

fitted results.

65

Chapter 4

Time resolved spin precession in Fe/CoO heterostructure

Optical excitation of spin precession is investigated in Fe/CoO (FM-

AFM) exchange coupled heterostructure with TRMOKE technique. Photo­

excited charge transfer processes in AFM CoO layer create a strong transient

exchange torque on FM Fe layer through FM-AFM exchange coupling, where

the efficiency of spin precession excitation is significantly higher and the

recovery is notably faster than the demagnetization procedure. The

precession amplitude peaks around room temperature and with external

magnetic field competitive to the magnetic anisotropy field, indicating that this

efficient excitation mechanism originates from the modulation of the uniaxial

magnetic anisotropy Ku induced by the FM-AFM exchange coupling. Our

results will help promote the development of low energy consumption

magnetic device concepts for fast spin manipulation at room temperature.

66

4.1 Introduction

Control of coherent spin precession in ferromagnets is currently a

popular topic due to its importance in magnetic recording and spintronic

devices [83-92]. Recently, optical excitation of coherent spin precession

utilizing ultrafast laser pulses provides novel approaches for fast spin

manipulation [83-89] and generation of ultrafast spin currents [86]. The

precessions are usually triggered by the transient torque on magnetization

generated via spin transfer effect [13, 90], demagnetization and modulation of

magnetic anisotropy [93-95]. The rapid demagnetization may be caused by

dissipation of spin angular momentum through electron-lattice interaction [83,

84], or transport of super-diffusive spin-polarized current [85], whereas the

anisotropy constant modulation may be caused by the alteration of spin-orbit

interactions due to hot electrons and electron-phonon thermalization [93],

creation of itinerant carriers [94], or coherent acoustic phonons [95].

Despite the different excitation mechanisms, most scenarios utilize the

optical absorption followed by a rapid raise in temperature and hence require

intense laser pulses. The long cooling time and excessive heating

considerably limit the push for fast and energy-efficient spin manipulation. For

instance, the demagnetization procedure usually requires intense optical

pluses to generate the non-equilibrium hot electrons [84] and thermodynamic

gradients [86]. In exchange-biased magnetic structures, the instant and

intense laser heating modifies the exchange bias, destroys pinning of the FM

spins at the interface, and thereby triggers the rotation of the magnetization to

67

the new direction [96]. But still, the thermal reduction of exchange bias is

followed by a few-hundred-picosecond recovery through heat diffusion.

The search for non-thermal excitation mechanisms motivates

extensive research to overcome the disadvantages of thermal ones [87, 88,

97-101]. The main idea is to utilize the interaction between magnetization and

photo-excited carriers that are selectively optical-pumped, where the

recombination time of photo-carriers is much shorter than the heat diffusion

process. For instance, optical spin-transfer [87] and spin-orbit [88] torques

are observed recently in FM semiconducting GaMnAs, which however usually

happens at low temperature. An alternative approach is through FM-AFM

exchange coupling, as small modulation of the exchange coupling strength

might lead to notable changes in magnetic properties [102, 103]. It is recently

demonstrated that short laser pulses can non-thermally introduce spin

reorientation and dynamics in AFM materials much faster than in FM

materials [104, 105]. But the question is still open whether it is possible to

drive FM at the speed of AFM through FM-AFM exchange across

heterostructure interface.

68

4.2 Samples and Characterizations

The Fe/CoO thin films are grown on the MgO(OOI) substrate in an

ultrahigh vacuum chamber. The MgO(OOI) substrate with a miscut angle less

than 0.5° is prepared by annealing at 600 °C for 30 min. The CoO thin films

are grown by a reactive deposition of Co under an oxygen pressure of

2 * 10-6 Torr. The thickness of CoO layer is 3 nm. Then, a 4 nm Fe film is

epitaxially grown on top of the CoO film at room temperature (RT). All of the

samples are covered by a 3-nm-thick MgO protection layer. The high quality

epitaxial growth of CoO film and Fe film can be proven by the reflection high

energy electron diffraction (RHEED) patterns [102], and the epitaxial relation

is CoO[110]//Fe[100],

Longitudinal MOKE measurements indicate a negligible exchange bias

(/-/e) in Fe/CoO from 80 K to above RT, as shown in Figs. 4.1 and 4.2. The

absence of exchange bias is likely due to a very small number of

uncompensated AFM spins at the smooth interface. However, the field

cooling leads to the alignment of AFM spins along CoO [110] directions,

which is collinear to Fe [100]. Due to the exchange coupling of the Fe

magnetization with the AFM spins, a uniaxial anisotropy appears below 7n

(discussed below), and the coercivity (Hc) increases with decreasing

temperature down to 130 K, below which Hc decreases because less AFM

spins are dragged to switch with FM spins [103]. As shown in Fig. 4.1, the

hysteresis loop is almost square for the field along the easy axis and is much

sharper at 80 K than 330 K; however, it is hard to reach saturation

69

magnetization at the largest field, restricted by our electromagnet, along the

hard axis perpendicular to the cooling field. Such results indicate a well-

defined easy axis and homogeneous anisotropy of the sample.

C Om mmm

c

■

-Qco

coco>

0 )

CD*

Figure 4.1 Longitudinal hysteresis loops for magnetic field along uniaxial easy axis

[100] and hard axis [010] at 82 K and 330 K.

8 2 K

H / / [ 1 0 0 ]

3 3 0 K

H / / [ 1 0 0 ]

- 1.0 - 0.5 0.0 0.5 1.0 Field(KOe)

70

600

^ 4 0 0

0

£ 2 0 0o 6

o

1 o

T -------------- 1-------------- 1-------------- ,---------------1-

HeHe

90 180 270 360

T(K)

Figure 4 .2 Temperature dependence of coercivity (Hc) and exchange biasing field

(He) obtained from easy axis loops.

71

4.3 Magnetic anisotropy

Figure 4,3 presents the TRMOKE results obtained from Fe/CoO thin

films with different magnetic field H along the Fe [110] direction (in-plane hard

axis) at RT. The oscillation is fitted to a damped sine function,

Aexp(-f/r )cos(2-nft), with precession amplitude A, time delay t , decay rate

1/r , and precession frequency f . It is obvious that spin precession frequency

f and relaxation time r increases as H increases. Figure 4.4 displays f as a

function of H. By solving the LLG equation, the frequency f can be well fitted

with the following equation:

2 = y J [ H cos( S - f ) + H a ] [ H cos(<?- j ) + H p] , (4.1)

where

Ha =Hdeff- 2KuS'm2flM s + K , ( 2 - sin22 <j>)IMs,

H? = 2K,cos40Ms + 2KuCOs2flMs, (4.2)

with effective demagnetization field Hdeff=4nMs+2K±/Ms, saturated

magnetization Ms, and gyromagnetic ratio y. </> and 5 are the angles of in-plane

equilibrium magnetization and H with respect to the Fe axis [100]. Ku, Ki and

Kx are the in-plane uniaxial, crystalline cubic, and out-of-plane magnetic

anisotropies, respectively. The <f> is derived from searching the minimum of

magnetic free energy. Therefore, the values of magnetic anisotropies at RT

are derived as: Hdeff= 2.0 kOe, KJMS=205 Oe and K1IMS=290 Oe.

72

•— 2112■ 1848■ 1320— 658 ► 461■ 330

264• 198— 132 >— 66

V

* *

05

— minus132

0.0 0.1 0.2- 0.1 0.3 0.4 0.5Delay time t (ns)

Figure 4.3 TRM OKE results of Fe/CoO thin films with different magnetic field H along

the Fe [110] direction (in-plane hard axis) at RT.

73

20

NXo

1000 2000 3000-1000 0

H (O e)

Figure 4 .4 spin precession frequency fa s a function of magnetic field H.

74

In order to study the uniaxial magnetic anisotropy Ku induced by FM-

AFM exchange coupling between Fe and CoO layer, temperature dependent

TRMOKE measurements are carried out. Figures 4.5 and 4.6 display the

TRMOKE results with magnetic field H= 3 kOe and 1 kOe applied along Fe

[11 yfl ] direction. The magnetization precession is faster at lower temperature,

due to the enhanced effective magnetic field Heff by the increase of Ku. Figure

4.7 summarizes the precession frequency f as a function of temperature.

Since the geometry of applied magnetic field is changed, the dispersion

relationship between f and H deduced from LLG equation requires

modification. By using the method mentioned in Chapter 2.2, we can deduce

the torque equations as follows:

\~dLt~ = = m<p [ ^ 1(2sin40sin2< cos2^ + 2sin20cos20 — sin20(sin4<p +

cos4<p) + 4sin20sin2<»cos2<p — cos20) + — (sin20sin2<p — cos2<p) +Ms

2cos26 + 2nMs) — ^ (sinBcoscp + sinQsin<p + a/2c o s 0)] ,

idmjp _ _ m&\^ (2 s m 2Bsin2(pcos2(pcos20 + 4sin29cos29sin2<pcos2(p +Y a t M s

cos40) + ^ Lsin2<pcos26 — 2cos29 + 2nMs j +^{sin9cos(p + sinBsincp +

yftcosB)] (4.3)

where 9 (<p) the angle between magnetization and Fe [001] (its projection on

(001) plane and Fe [100]) direction in the spherical coordinate system. By

treating the mg and with wave form as in Chapter 2.2, one can derive:

2kf = yJW Jh (4.4)

75

3nj,lccg»to

a)*ca)'</>cCO

j * ‘ . * •* *''** .*—*7 \ V a ^ - ^ wr \

»*S4

A U * U * ~ ^s .-. •. ; * ♦/« y ***

* * 4 « 4 4 ^ ^ _

-t._/\ yv y \ v \ - / ,%-

300K290K280K270K260K250K240K230K210K190K150K100K50K10K

0.0 0.1 0.2 0.3 0.4 0.5

t(n s )

Figure 4.5 Temperature dependent TRM OKE results with magnetic field H= 3 kOe.

76

Tran

sien

t Ke

rr si

gnal

(a.

u.)

• • ■. / ......

300K290K280K270K260K250K240K230K210K190K150K100K50K10K

Figure 4 .6 Temperature dependent TRM OKE results with magnetic field H= 1 kOe.

77

35

30

_ 25NXo

20

15

Figure 4.7 spin precession frequency f as a function of temperature T.

«

H = 1 kOe H = 3 kOe

0 50 100 150 200 250 300

T ( K)

78

The next step is to determine the equilibrium direction of magnetization,

i.e. the values of 0 and (p. The magnetic free energy F can be written as:

F = — (sin4 0sin2 2d> + sin2 20) + K u sin2 ^sin2 0 + K, cos2 0 + 2 /r(M cos#)24 (4.5)

- (sin 0 cos <j> + sin 0 sin <f> + 42 cos 0)

Along the equilibrium direction of magnetization, F reaches the

d F dFminimum, so that — = 0 and — = 0. Therefore, 0 and q> can be expressed

as a function of anisotropy fields and H. By fitting the field dependence of f,

one can obtain the equilibrium direction of magnetization as well as the

anisotropy fields.

f 25

O.

15 -

100 150Temperature (K)

200 250 300

Figure 4.8 q> as a function of temperature (green=3 kOe, blue=1 kOe).

79

thet

a (d

egre

e)

88.5

87.5

86.5

85.5

84.5

100 150Temperature (K)

200 250 300

Figure 4.9 G as a function of temperature (green=3 kOe, blue=1 kOe).

80

2.0

0O

Ku/Ms

0.5

0.00 100 200 300

T (K)

Figure 4.10 Uniaxial magnetic anisotropy field KJM S as a function of temperature T.

Figure 4.10 shows the dependence of uniaxial magnetic anisotropy

field KJMs on the temperature. Ku increases as T decreases. This indicates

that its origin is the AFM ordering of CoO layer. The FM-AFM exchange

coupling between Fe and CoO layer establishes an extra preference of

magnetization alignment in Fe, where the FM spins favor aligning parallel with

the frozen AFM spins. As the number of frozen AFM spins in CoO is

increased with lower T, the FM/-AFM exchange coupling is enhanced, which

leads to an increase of Ku.

81

4.4 Ultrafast spin exchange torque via photoexcited charge

transfer processes

TRMOKE measurements are performed in a pump-probe setup, where

the intensity ratio of the pump to probe pulses is set to be about 6:1. The

probe (A=800 nm) utilizes the MOKE technique with crossed polarizers and

incident angle around 40° to investigate the transient magnetic state change

along longitudinal and polar directions. Two strategies of pump are employed

in the measurements to investigate the optical excitation mechanisms: First,

the more intense pump pulses (A=800 nm, 3.1 mJ/cm2) are used to modulate

the FM order of Fe layer as shown in Fig. 4.11 (left), since the CoO layer is

almost transparent to 800-nm light [106]; Second, the weaker pump pulses

(A=400 nm, 0.16 mJ/cm2) are utilized to mainly affect the AFM order of CoO

as depicted in Fig. 4.11 (right). All TRMOKE measurements are performed

after field cooling the sample.

The TRMOKE result from Fe/MgO heterostructure is displayed in Fig.

4.12 (black squares) with pump pulse fluence 3.1 mJ/cm2 and magnetic field

H= 2 kOe at room temperature. The sudden rise and decay of Kerr signal

corresponds to the demagnetization procedure as described by the three-

temperature model [107, 108]. The rise in electron temperature by

demagnetization is first partially relaxed to the lattice within several

picoseconds. Afterwards, the heat of electron, spin and lattice diffuses to the

surrounding over >500 ps, which slowly recovers the system to its ground

82

state. Meanwhile, the magnetization starts to precess back to the previous

equilibrium direction in a damped circling way described by LLG equation. As

a result, the measured Kerr signal can be well fitted by the following equation

6|<= a*exp(-t/?o)+A*exp(-f/ r)sin(2ir ft+q>), (4.6)

where parameters A, r, f and <p are the amplitude, magnetic relaxation time,

frequency, and initial phase of the magnetization precession mainly along the

polar direction, respectively. Here, a and to are related to the background

signal owing to the slow recovery of magnetization after fast demagnetization

by the pump pulses, which happens mainly along the longitudinal direction.

The amplitude a qualitatively describes how far the FM system is driven away

from the ground state, and the amplitude ratio A/a can be used to qualitatively

estimate the efficiency of optically-driven coherent spin precession, which is

around 0.16 for Fe/MgO (black squares in Fig. 4.12).

Figure 4.12 (red circles) shows the TRMOKE result from Fe/CoO(3

nm)/MgO structure with the same excitation condition as the measurement on

Fe/MgO (black squares). We note that the background amplitude a remains

unchanged between the two measurements, as shown in Fig. 4.12, while in

contrast the amplitude of coherent spin precession A is enhanced with A/a =

0.85 for Fe/CoO/MgO heterostructure (red circles). This shows that the AFM

CoO layer improves the efficiency of optically excited coherent spin

precession. In Fe/MgO, A can be understood as Ai+Am, where A and Am

denote the contributions from modulation of magnetic anisotropy field Ha

83

through reduction of cubic magnetic anisotropy constant Ki and

magnetization Ms, respectively. Whereas in Fe/CoO, there exists another

contribution Au, which originates from modulation of the uniaxial magnetic

anisotropy constant Ku induced by FM-AFM exchange coupling between Fe

and CoO layers. Since Ku is more sensitive to temperature rise compared with

Ki [103], the modulation of Ku through laser pulse heating contributes

significantly to the enhancement of spin precession excitation.

The key finding here is that pronounced spin precession is still

observed with much lower pump pulse energy 0.16 mJ/cm2 at 400-nm

wavelength, as shown in Fig. 4.12 (blue triangles). Moreover, the TRMOKE

data reveal the absence of obvious demagnetization and slow recovery of

a Ms. Therefore, the rise of electron temperature Te in Fe layer through laser

pulse heating is negligible by the application of low pump pulse intensity.

Thus, the amplitude ratio A/a =8.35 is about two orders of magnitude larger

than from Fe/MgO with stronger pump pulse. Also, the instant pronounced

spin precession points towards an ultrafast non-thermal excitation process in

the AFM CoO layer as depicted in Fig. 4.11 (right) and discussed further

below. Furthermore, the magnetic relaxation time t decreases from 330 ps

(red circles) to 100 ps (blue triangles) in Fe/CoO with gentler pump pulses,

which is desirable for fast switching [109]. This is due to the fact that Gilbert

damping is enhanced at lower temperature in Fe/CoO exchange system [103].

84

800 nm 400 nm

®@®@@ ® @ ®

XX• • • •

FM Fe -

FM/AFM exchange

— AFM CoO —

*@@@@

i Co 3d

3 > 0 2 ph v ~ 3 . l e V

N(E)

Figure 4.11 Two different pump strategies to optically excite the coherent spin

precession, where the black arrows represent the spins of magnetic ions. The optical

charge transfer transition in CoO is depicted.

85

=3aj

roca'tok.fc

*

0 200 400Delay t(ps)

Figure 4.12 TRM OKE results from Fe/MgO (black squares) and Fe/CoO (red circles)

with pump pulse intensity 3.1 mJ/cm2, and Fe/CoO (blue triangles) with pump

intensity 0.16 mJ/cm2.

- Fe/MgO 3.1 mJ/cm2 • Fe/CoO 3.1 mJ/cm2 a Fe/CoO 0.16 mJ/cm2

• \ j l ,A=800nm

^ 'A A A A a aA=400 nm

86

The optical excitation of spin precession in Fe/CoO heterostructure can

be due to either the demagnetization procedure, or a modulation of magnetic

anisotropy constants, or spin transfer torque, or a combination thereof. From

the results described above, the fast demagnetization procedure is not

observed and the saturated magnetization Ms remains almost unchanged,

where AM scales down with lower a. Moreover, the modulation of magnetic

anisotropy constant Ki in Fe layer is reduced by the decrease of pump pulse

energy (i.e. Ai also contributes much less). Another possibility of the

enhanced Ala is the spin transfer induced torque. Even though the CoO

insulating layer eliminates the spin current transfer via spin pumping effect [9],

we have shown recently that the FM/AFM exchange establishes a spin

angular momentum transfer channel in Fe/CoO heterostructure, which slightly

increases with lower temperature [103]. To check this possibility, temperature

dependent TRMOKE measurements are carried out. Figures 4.5 and 4.6

present the TRMOKE results with pump pulse intensity 0.16 mJ/cm2 at

different temperatures T, where the absence of demagnetization peak is

observed ranging from low to room temperature. Figure 4.13 summarizes the

precession amplitude A as a function of temperature T with different magnetic

field H applied, and shows that A is peaked around 270K-290K, indicating

that spin angular momentum transfer is not the main contribution. Therefore,

the origin of this efficient excitation mechanism has to be the modulation of

magnetic anisotropy constant Ku as a result of variance in FM-AFM exchange

coupling.

87

H = 0.7 kOe H = 2 kOe20

< 10

01 0 0 200

T (K )300

Figure 4 .13 Measured and simulated spin precession amplitude A as a function of T.

88

T= 240K

T= 300K

T= 78K

00 1 2 3

H (kOe)

Figure 4.14 The spin precession amplitude A is plotted as a function of H at different

T, where solid lines represent simulated results.

89

The modulation of Ku also sheds light on the difference between the

two optical excitation strategies, 800-nm versus 400-nm wavelength pulses.

On one hand, the CoO layer is almost transparent to the light with A=800 nm,

as its bandgap is around 2.5 eV [106]. Therefore, the 800-nm pump pulses

mainly raise the electron temperature in FM Fe layer as depicted in Fig.

4.11 (left), and modulate the FM-AFM exchange coupling and hence Ku. This

leads to improvement of spin precession excitation efficiency as compared

with Fe/MgO thin film (Fig. 4.12 black squares and red circles). It is also

reported by Ju. et al. [96] that the heat in FM layer partially diffuses to the

adjunct AFM layer and slightly modulates the interface AFM order. On the

other hand, the FM order in Fe layer is much less modulated by the gentle

pump pulses with A=400 nm (3.1 eV), while there exist photon-excited

electron transitions in CoO layer as depicted in Fig. 4.11 (right). The near-gap

photons excite charge transfer transitions from O 2p band to Co 3d band that

is partially occupied with minority spins [110], which promotes the nearest-

neighbor FM exchange interaction and reduces the AFM ordering in CoO

[111], Since the coupling between Fe and CoO is a long-range interaction

[112], the variance of AFM order in CoO layer notably modulates the FM/AFM

exchange coupling and Ku, and thus leads to significant enhancement of the

excited spin precession (Fig. 4.12 blue triangles).

In order to derive the value of Ku, magnetic field dependent TRMOKE

measurements are carried out at different T. Figure 2(c) (black squares and

red circles) displays the spin precession frequency f for two different fields H

90

as a function of 7, where f increases as H increases. This can be understood

by the enhancement of the effective field due to larger H and is well fitted with

LLG equation to derive the magnetic anisotropies. It turns out that the cubic

magnetic anisotropy Ki originating from Fe layer remains almost unchanged,

while the uniaxial magnetic anisotropy Ku induced by the FM/AFM exchange

coupling between Fe and CoO layers increases significantly with decreasing

T, as shown by the blue triangles in Fig. 4.10. The derived KJMs behaves

similarly like f as a function of 7. The FM-AFM exchange coupling establishes

an extra preference of magnetization alignment in Fe, where the FM spins

favor aligning parallel with the frozen AFM spins. As the number of frozen

AFM spins in CoO is increased with lower T, the FM/AFM exchange coupling

is enhanced, which leads to an increase of Ku and hence increases f with

lower 7 as shown in Figs. 4.7 and 4.10.

To further elucidate the coherent spin precession excited by

modulation of anisotropy constant Ku, the relationship between precession

amplitude A and applied field H is investigated. Although the modulation of

anisotropy field Ha is independent of H, the transient torque on the

magnetization becomes stronger with increasing H as it pulls the

magnetization away from Ha, thus leading to larger precession angles at

higher fields. When the magnetization is nearly aligned along H at the field

strength stronger than Ha, the torque starts to saturate and the precession

angle decreases with increasing H. Thus, one can expect that the precession

amplitude reaches a maximum at the field approximately equal to the static

91

anisotropy field. Figure 4.14 shows the value of A as a function of H. We note

that at T = 300K the amplitude A (black squares) is first sharply enhanced

with increasing field and then saturates around H = 1 kOe and at last slowly

decreases, consistent with our expectation. Moreover, the simulation of A

versus H is carried out, where A is assumed to be proportional to the

equilibrium direction change by modulation of Ku (AKU) [113-115]. The AKU

and scaling factor s are parameterized at 300K. The black curve in Fig. 2(d)

corresponds to the simulation of A with a reduction of Ku about (165 Oe*Ms),

which agrees quite well with the measured relationship between A and H.

The simulation results also shed light on the evolution of precession

amplitude A with temperature T shown in Fig. 4.13. As T decreases, Ha is

enhanced by Ku, which requires stronger H to compete with Ha and the

generated transient torque is saturated. As a result, the peak in amplitude

peak is shifted to higher fields and becomes broader with lower T (Fig. 4.14).

The blue curve in Fig. 4.14 presents the simulated result at T=240K with s as

derived above and AKU parameterized equivalent to the effect of 20-K rise in 7.

The amplitude peak is around H - 1.8 kOe and the rise and decay of A due to

increasing H is less steep than at 7=300K. The shifting and broadening of the

amplitude peak is further enhanced at 78K as illustrated by the measured

results (red circles). Therefore, the evolution of amplitude A with temperature

7 in Fig. 4.13 can be interpreted as the shift of amplitude peak. Figure 4.13

shows the simulated amplitude A as a function of 7 for two different fields H

(dashed lines), which agrees quite well with the measured values. In this

92

simulation, AKU corresponds to a temperature rise of A7=20K, which is much

larger than AT «1K in CoO layer due to direct heating from optical absorption.

The photo-excited charge transfer processes essentially raise the

temperature of the CoO spin system, which effectively modulates the AFM

order of CoO.

In order to investigate the duration of this fast FM/AFM exchange

torque, simulations of real-time magnetization precession are carried out from

LLG equation. The observed pronounced coherent spin precession starts

right at t= 0, indicating a sudden change of equilibrium direction caused by

AKu. W e assume that the reduction of Ku (165 Oe*Ms) happens instantly at

t=0, followed by an exponential recovery process with time constant rr. Real

time simulation of magnetization precession is carried out with LLG equations

with time interval Af=0.2ps:

mz{t+At)IMs= mz(t)/Ms*exp(1-At/T) - my(t)l Ms *Y (Hcos(6-(p)+H^(t))*At

my(t+At)/Ms= my(t)/Ms*exp(1-Af/T) + mz(f+Af)*y (Hcos{6-<p)+hf(t))*At +(<p(Ku -

AKu(t+At))- <p(Ku -AKu(t))) (4.7)

where the first term represents the Gilbert damping effect and the second

term denotes the circling of spin precession. As for the my, there is an

additional term describing the change of equilibrium direction driven by

modulation of Ku.

As shown in Figs. 4.15 with rr = 40 ps, the simulated time evolution of

magnetization (red curve) agrees quite well with the probed Kerr signal (black

93

dots), since the TRMOKE signal is proportional to the polar component of

magnetization precession. Shorter or longer t> leads to mismatch in the

oscillation phase, as shown in Figs. 4.16-4.18. The fast recovery from photo­

excited transitions might be due to the strong carrier-phonon interaction,

which results in non-radiative and phonon-assisted carrier recombination

[116]. Moreover, the recovery time of the instant photo-induced exchange

torque is much faster than the cooling time from demagnetization, which

might promote novel device concepts for fast spin manipulation.

step size = 0.8 ps xr = 40 ps £0.02

0.01

s 0.00o>

- 0.01AAT,

- 0.02

800 40Delay t (ps)

Figure 4.15 Measured Kerr signal change with 0.8 ps step size (black squares),

simulated polar magnetization component (red curve) and assumed modulation of

Ku (blue curves) as a function of t.

94

0.02

0.01CO

§ 0.00 N E

- 0.01

- 0.02

0 200 400Delay t (ps)

Figure 4.16 Measured Kerr signal change with 4 ps step size (black squares),

simulated polar magnetization component (red curve) and assumed modulation of

Ku (blue curves) as a function of t.

step size = 4 ps xr = 40 ps

305*

"rocg>co

a>

95

step size = 0.8 ps Tr = 2 ps ^0.001

ni 0.000

- 0.001

0 40Delay t (ps)

Figure 4.17 Measured Kerr signal change (black squares) and simulated polar

magnetization component (red curve) with rr = 2 ps.

96

Kerr

sign

al (

a.u.

)

0.02step size = 0.8 ps Tr = 15 ps / y *

0.01

* 0 . 0 0

- 0.01

0 40 80Delay t (ps)

Figure 4.18 Measured Kerr signal change (black squares) and simulated polar

magnetization component (red curve) with rr = 15 ps.

97

Kerr

sign

al (

a.u.

)

0.02

0.01CO^ 0.00

£

- 0.01

- 0.020 40 80

Delay t (ps)

Figure 4.19 Measured Kerr signal change (black squares) and simulated polar

magnetization component (red curve) with rr = 100 ps.

_ - step size = 0.8 ps — xr = 100 ps

98

Kerr

sign

al (

a.u.

)

Chapter 5

Spin precession excitation in manganite thin films

Pronounced spin precessions are observed in a geometry with

negligible canting of the magnetization in FM Lao.6 7Ca0 3 3 Mn0 3 thin films using

the TRMOKE technique. The precession amplitude monotonically decreases

with increasing field, indicating that the coherent spin rotation may be

triggered by a transient exchange field and not by demagnetization and/or

anisotropy field modulation. W e attribute the transient exchange field to

emergent AFM interactions due to charge transfer and modification of the

kinetic energy of eg electrons under optical excitation.

99

5.1 Introduction

Control of coherent spin rotation in ferromagnets has received intense

investigation in the past decade for fast magnetization switching in novel

magnetic recording and spintronic devices [91,117], Recently, optical

excitations of coherent spin precessions have been demonstrated in various

FM metals, providing a new approach for fast spin manipulation. The

precessions are triggered by effective field pulses generated via fast

demagnetization [118,119], or transient modulation of anisotropy constants

[93-95], The demagnetization occurs due to fast dissipation of spin angular

momentum via efficient spin-lattice thermalization [83], whereas the

anisotropy constant modulation may be caused by the alteration of spin-orbit

interactions due to hot electrons and electron phonon thermalization [93],

creation of itinerant carriers [94], or coherent acoustic phonons [95],

Despite the different excitation mechanisms, both scenarios require

canting of the magnetization, i.e., an angle between magnetization and the

external applied magnetic field to induce the coherent rotation. This geometry

is also essential for spin precession in exchange biased magnetic structures

[96], In such systems, the instant laser heating destroys the exchange

induced pinning of the FM spins at the interface, and thereby triggers the

rotation of the magnetization to the new direction. Djordjevic et al. recently

proposed a new excitation mechanism, where localized spin flip may decay

into short wavelength spin wave excitations and then transfer to low energy

spin wave modes [120], The canted geometry also may not be important for

100

ferrimagnetic materials in which the spin precession is excited via inverse

Faraday effect using circular polarized light [104,121].

In this Chapter, observations of pronounced spin precessions are

reported in a geometry with negligible canting of the magnetization in FM

La06 7Ca0.3 3MnO3 (LCMO) thin films on NdGa03 (NGO) substrates. Using

TRMOKE measurements, large spin precessions are observed over a wide

range of magnetic fields (0 -2 .5 T) applied along the in-plane easy axis. The

precession amplitude monotonically decreases with increasing field, in

contrast to the precession induced by the transient anisotropy field in the

state of canted magnetization. These observations indicate that the coherent

spin rotation may be triggered by a transient exchange interaction and not by

demagnetization and/or anisotropy field modulation. The picosecond

modulation of the exchange interaction results from the emergence of AFM

interaction caused by charge transfer and alteration of kinetic energy of eg

electrons in LCMO under optical excitation.

101

5.2 Samples and Experiments

The LCMO films are epitaxially grown by pulsed-laser deposition on

NGO or SrTi0 3 (STO) substrates. An excimer laser with a wavelength of 248

nm, an energy density of 2 J/cm2, and a repetition rate of 5 Hz is used. The

growth rate is around 0.04 nm/pulse, and the thickness of the film was

determined by the nominal value. The substrate temperature is around 780

°C and the oxygen pressure was about 500 mTorr during deposition. The

films are cooled down in about 400 Torr of oxygen at the rate of 20 °C/min

after deposition. Unstrained films were obtained by annealing the films in a

quartz tube at 900 °C with a flow of 50 standard cm3/min pure O2 for 24 h and

cooled at the rate of 10 °C/min to room temperature. The films on NGO

substrates exhibit an inplane uniaxial anisotropy of 0.2 T at thickness of 6 0 -

100 nm [122],

TRMOKE measurements are performed with a Ti:sapphire amplifier

laser system providing 150-fs pulses at 1-kHz repetition rate. W e use 800-nm

pump pulses with a fluence of 1 mJ/cm2 to excite the magnetization in our

samples, and the time evolution of the magnetization is monitored with 400-

nm probe pulses, where the Kerr rotation is at a maximum [123], with

fluencies on the order of 0.1 mJ/cm2. The transient Kerr rotation Ad is

measured in a cross polarization configuration. Rotation of the analyzer

through the extinction angle leads to the opposite sign of Ad. For transient

reflectivity ( AR ) measurements we use the same pump and probe

wavelengths as in TR-MOKE.

102

5.3 Coherent Spin Precession via Photoinduced Anti­

ferromagnetic Interactions in Lao.67Ca0.33Mn03

Figure 5.1(a) shows the time evolution of Ai? from a 100 nm thick

LCMO film at 25 K. AR exhibits a fast response with a rise time comparable to

the laser pulse width and a slow response with decay time of 100 ps. The

latter component is associated with the slow spin-lattice relaxation [124].

The time evolution of Ad is shown in Fig. 5.1(b). Here, the external field

H is applied perpendicular to the sample plane and the probe beam is

incident at an angle of 5 degrees with respect to the sample normal,

permitting sensitive detection of the out-of-plane magnetization component

Mz. In contrast to the fast electronic response, the magnetic response shows

no instantaneous change on the time scale of 1 ps. The major feature of the

magnetic response manifests as strong oscillations which correspond to the

coherent precessions of the FM spins. W e therefore conclude that no

significant demagnetization occurs before the magnetization rotation is

launched. This finding is consistent with the observations of negligible

demagnetization in FM Lao.6Sr0.4Mn0 3 films at temperatures significantly

lower than its Curie temperature [123], The weak demagnetization is due to

the half metallic nature in manganites where the Elliot-Yafet scattering is

blocked [107],

103

3m

5

0 50 100 150

H I I Easy Axis (c)

Time (ps)

H x Plane (b)

3

■<3

0 200100

0.05 T

0.3 T

■ V

Time (ps)50 100 150 200

Time (ps)

3

% H = 0 T

0 200 400 600 0Time (ps) Frequency (GHz)

Figure 5.1 Transient reflectivity AR (a), and transient Kerr rotation Ad of LCMO film

with magnetic field H (0.5 T) applied perpendicular to the sample plane (b), with H

applied along the in-plane easy axis (c), and with zero field (d) at 20 K; Fourier

transform (e) of Fig. 5.1(d).

104

easy axis (b)

560 Oe

OOe

200 4000

hard axis (a)

« 560 Oe

200 4000Tim e (ps) Tim e (ps)

Figure 5.2 Transient Kerr rotations of Fe film with magnetic field applied along the in­

plane hard axis (a) and easy axis (b) at RT.

105

Ferromagnetic metal Half metal

Figure 5.3 Three-temperature model in ferromagnetic metals and half-metals.

As a comparison, Fig. 5.2(a) shows the laser induced magnetization

dynamics in an Fe film at an external field of 550 Oe applied along the in­

plane hard axis at room temperature (RT). The transient Kerr rotation clearly

shows a fast response which occurs within 1 ps. This response originates

from the fast demagnetization due to the spin flip process associated with the

spin-lattice interactions [118]. The demagnetization decays within 50 ps, while

the subsequent magnetization precessions persist for much longer time (>

500 ps) [125]. The excitation of the magnetization precessions in the Fe film

can be reasonably explained by the modulation of the anisotropy fields

caused by the demagnetization [118,119]. Under the laser interaction, the fast

demagnetization modifies the magnetocrystalline anisotropy, generating a

transient effective field Htr perpendicular to the static effective field Hen, and

thus the magnetization precessions are triggered.

Hence, the precession is significantly weaker and difficult to observe if

the external field is applied nearly close to the easy axes along which the

magnetization is pinned, as shown in Fig. 5.2(b). These observations clearly

indicate that magnetization canting and fast demagnetization are of crucial

importance in exciting spin precession in the Fe film, typical for 3d transition

ferromagnetic metals. Figure 5.3 depicts the difference of three-temperature

model in ferromagnetic metals and half-metals.

In contrast, Fig. 5.1(b) shows that fast and large demagnetization is not

present in the spin precession dynamics of LCMO films. In the following, we

will present further measurements of the magnetization dynamics when H is

107

applied along the uniaxial easy axis in the film plane, i.e., a configuration with

negligible canting of magnetization. The results are shown in Fig. 5.1(c).

Similar to the case where the magnetization is pulled out of plane, no fast

demagnetization occurs when spins are aligned within the film plane, while

pronounced spin precession is present over a wide range of magnetic fields.

Particularly, a large precession is also excited at zero field and persists for

over 1 ns, as shown in Fig. 5.1(d). We note beating patterns in the transient

Kerr spectra at zero field. Fourier transform reveals two spin waves modes at

frequencies of 11 GHz and 15 GHz (Fig. 5.1(e)). The mode with lower

frequency corresponds to the Kittel mode, as the detailed analysis of the

magnetic field dependence of the precession frequency shows [122]. The

higher-frequency mode may correspond to a surface Damon-Eshbach-type

mode or a higher-order standing spin wave mode within the ferromagnetic

film [126,127], The slow decay of the precession indicates that the dephasing

process is weak, and thus indicates a well-defined easy axis along which all

exchange-coupled spins are aligned at zero external field. In this case canting

of magnetization does not exist, and therefore modification of the in-plane

anisotropy may not be the main contributor to the transient field that launches

the spin precession shown in Fig. 5.1(c).

In the state of canted magnetization, anisotropy modulation will indeed

generate a transient field, triggering the spin precession. The anisotropy

modulation may arise from the thermal modification of magnetocrystalline

energy at elevated electron temperature [93]. Although the modulation of

108

anisotropy field Ha is independent of the applied field H, the transient torque

on the magnetization becomes stronger with increasing H as it pulls the

magnetization away from Ha, thus leading to larger precession angles at

higher fields. When the magnetization is nearly aligned along H at the field

strength stronger than Ha, the torque starts to saturate and the precession

angle decreases with increasing H. Thus, one can expect that the precession

amplitude reaches a maximum at the field approximately equal to the static

anisotropy field. The top left panel of Fig. 5.4(a) shows the field dependence

of precession amplitude obtained when H is applied at an angle of 45 degree

to the easy axis, a state with canted magnetization. We note that the

amplitude is sharply enhanced with increasing field and then slowly reduced

at fields larger than 0.3 T, consistent with our expectation. Similar results are

obtained from the other LCMO samples, as well as the Fe film with the field

applied along the hard axis and the direction at an angle of 10 degree to the

easiest axis (Fig. 5.4(b)). The modulation of the magnetic in-plane anisotropy

in an oxide ferromagnetic material C r02 also leads to similar dependence of

precession amplitude on H applied along the hard axis, and no precession is

detected if H is oriented close to the easy axis [93,128].

In contrast, the precession with the applied field along the easy axis of

LCMO shows the opposite dependence; i.e., the amplitude decreases with an

increasing field for all three LCMO samples, as shown in the right panels of

Fig. 5.4(a). This result confirms our previous conclusion that the transient

anisotropy field due to demagnetization or modification of anisotropy constant

109

is not the major contributor to launch the spin precession when H is applied

parallel to the easy axis. W e thus propose a transient exchange field H lrx that

triggers the spin precession.

(a)2 ■

«n3

. cSrEPn .D a hard axis■ □ □

□,S 60 nm

0.9

0.6

0.3

1-O

■%9 0o

LCMOeasy axis

□□

X1 i

1

1.2 o

D □ • □ 0.8 o

^ 100nm0.4

o o

1.2

■ ■ □ □

□ 0.8o

o. □g 100 nm* □...1.... ■.....A-... .......L

0.4J— ......

oOO

.1..... -..... 1...1 2 0

Magnetic Field (T)

(b>

T>

O.E<

1 •

□ □ 6 £0oo01

□□

cs

■'—r □

1□ 4 o Fe

-o.q Iff to easy axis1 . 1 i 1

hard axis□I * ....A - 1

2

0 400 800 0 400 800Magnetic Field (Oe)

Figure 5.4 Amplitude of precession of 60 nm and 100-nm LCMO films with field

applied at 45 degree to the easy axis and along the easy axis at 20 K (a), and Fe film

with field applied at 55 degree to the easy axis and of 10 degree to the easy axis at

RT (b).

110

In order to obtain the field dependence of H lrx , we simulate the

precession of the magnetization and follow its out-of-plane component. In the

simulation, the magnitude of //|£ is adjusted so that the calculated spin

precession angle coincides with the nominal precession angle obtained from

the measured precession amplitude (Fig. 5.4(a)). W e thus determine the field

dependence of the magnitude of H lrx for the 60-nm thick film, as shown in Fig.

5.5(a). We note that Hlrx is reduced with increasing applied field. This

dependency is a characteristic of the transient exchange field in manganites

as discussed below. In contrast, the transient anisotropy field H£r , which is

perpendicular to the magnetization, due to demagnetization or reduction of

the anisotropy constants is enhanced with increasing field. We present a

calculated result in Fig. 5.5(b), which corresponds to t f ' a in a configuration

with slight canting of magnetization, i.e., an angle of 5 degrees between the

applied field H and the uniaxial anisotropy field Hu, and 10% reduction of Hu

(0.2 T) after laser interaction.

^ 10

50o>

Z 40

30

0 1 2 M agnetic R eid (T )

0 1 2 M agnetic Field (T )

Figure 5.5 Transient exchange field (a) and transient anisotropy field (b) in a state of

canted magnetization (magnetic field at 5 degree to the easy axis) of 60-nm thick

LCMO film. The dashed line is a guide to the eye of the data.

I l l

The transient exchange field may result from the modification of the

exchange coupling because the optical excitation may change the balance

between the FM metallic and AFM insulating states in the film. The

coexistence of distinct FM metallic and AFM insulating phases in doped

manganites has long been established [129-131], and the AFM domain may

be melted in the presence of an external field. In the case of LCMO, several

reports indicated the presence of the AFM phase down to the lowest

temperatures [132-134], Very slow photoinduced drops in conductivity and

demagnetization have been observed in LCMO and interpreted as the

photoinduced shift in the FM/AFM phase balance triggered by the

photoexcited carriers that promote the cooperative Jahn-Teller distortions and

formation of AFM insulating clusters [135]. An ultrafast drop in conductivity

upon photoexcitation with femtosecond pulses was observed by Averitt et al.,

which is also consistent with the promotion of AFM correlations [124],

We speculate a formation of nanometer- to sub-micrometer- scale

regions with transient AFM interaction after the intense optical excitation,

similar to the size of AFM domains observed in quasistatic phase separated

manganites [129-132], Since the AFM regions are randomly distributed in

LCMO films, we would expect a spatial inhomogeneity of the transient

exchange field which may lead to excitations of the Damon Eshbach mode or

the standing spin waves [127,128]. Because application of the magnetic field

may suppress the AFM interaction, the magnitude of Hlrx is reduced at higher

fields as shown in Fig. 5.5(a).

112

The reorientation of the spins to form AFM clusters may take a long

time due to the slow spin-lattice thermalization process, a general feature in

manganites [136], to dissipate spin angular momentum. However, the

exchange interaction may be modified in a much faster time regime. This is

because the local carrier concentration and its kinetic energy play a crucial

role in determining the magnetic phases [137,138], Both photoinduced charge

transfer between neighboring Mn ions [139,140] and the thermalized

electrons may suppress the double exchange interaction and promote the

AFM coupling. Therefore, the exchange field pulse may be generated within

electron phonon relaxation time, much faster than the emergence of the

demagnetization caused by the real formation of AFM clusters.

This finding is corroborated by the temperature dependence of the

magnetization dynamics shown in Fig. 5.6. At temperature lower than 150 K,

the magnetization dynamics show negligible demagnetization. With

increasing temperature, a slow demagnetization (> 500 ps) is clearly

observable. However, the spin precession is launched nearly at the same

time at all temperatures lower than the Curie temperature Tc (~270 K). The

reduction of the precession amplitude at higher temperatures is mainly

caused by the decrease of the saturated magnetization. The reduction of the

magnetization also results in a pronounced decrease of precession frequency

at temperatures approaching to Tc, as shown in the inset of Fig. 5.6. This

result therefore supports our model that the precession is triggered by the

transient exchange field in the absence of spin-lattice thermalization.

113

■

3a

• w

<J

' m T=150 K

J j a -

• • • - • i -

T=230 K

r '.

|V T=260 K1

A

i ....». i ...... ..... —i—

■ -.2 0 N5 10

T=2

.............. 'SHMX5 T

) 100 200 3(T (K)

90 K

i...- ■.....

)0

10 250 500 0 250 500

Tim e (ps)

Figure 5.6 Transient Kerr rotations of LCMO film with applied field H=0.5T at various

temperatures. The inset shows the uniform precession frequency as a function of the

temperature.

114

A recent study reveals a transient anisotropy field due to the optical

induced transition between easy in-plane and easy out-of-plane anisotropy

states in La2-2xSri+2xMn20 7 bilayered manganites [94], However, we point out

that the LCMO film is not located at the boundary of two ferromagnetic

phases with different anisotropies, and therefore no transient field due to the

emerging of anisotropy field along different directions is expected to be

present in our LCMO films. Such a transient field also would be independent

on the external field, different from H lrx which we obtain from the simulation of

the field dependence of the precession amplitude as shown in Fig. 5.5(a).

115

Chapter 6

Magneto-electric coupling at PbZr0.52Ti o.4803/Lao.67Sro.33Mn03

interface

The interfacial spin state of the multiferroic heterostructure PbZr052 Ti

o.4803/Lao.67Sr0 3 3 Mn0 3 (PZT/LSMO) and its dependence on FE polarization is

investigated with MSHG technique at 78K. The spin alignment of Mn ions in

the first unit cell layer at the heterointerface can be tuned from FM to AFM

exchange coupled, while the bulk magnetization remains unchanged. Multiple

domains of both phases coexist as the ferroelectric polarization is switched.

The results will help promote the development of new interface-based

functionalities and device concepts.

116

6.1 Introduction

For the past few decades, magnetic recording has been a major

technology for information storage, while binary information is stored as

magnetization of magnetic elements. During this period of miniaturization of

magnetic memory cells, the reading procedure turns out to be scalable down

to the nanometer range. However, the traditional writing mechanism driven by

a magnetic field appears as a crucial limiting factor, due to the large power

consumption and loss of accuracy caused by the magnetic stray fields

affecting neighboring cells. In the modern MRAMs, spin-polarized current

induced spin-torque is used to switch magnetization states instead of

magnetic field, which somehow enables further down scaling of unit cells [9-

12]. Nevertheless, the large current densities required also lead to a

significant energy loss from heating. Thus in the push for low-energy

consumption logic devices, electric field control of magnetization switching via

strong magneto-electric coupling seems to offer a new mechanism to break

through those limitations and improve device performance; as such, it has

recently drawn increasing research interest.

Transition metal oxides are promising candidates for multifunctional

devices, due to their strong correlated electron system with coupled charge

[141], spin [142] and orbital [143] degrees of freedom. This route is carried

out by fabricating well-defined interfaces between these complex oxides, to

engineer and cross-couple their unique electric, magnetic, and transport

properties [144, 145]. The interfaces of TMO heterostructures have the

117

genuine property of breaking space inversion symmetry, and thus offer a

unique and important test-bed to promote new phases and properties that can

be controlled at atomic precision, such as conducting electron gases [141],

superconductivity [146] or electric polarization dependent spin transfer

[15,147],

The interfacial spin configuration of oxide heterostructures under

electronic and structural reconstruction is key to emerging multiferroic and

spintronic technologies with new functionality, and thus attracts research with

different approaches [148-152], However, most experimental probes are

indirect due to the lack of sufficient resolution to locate spin configuration at

the atomic scale. The interfacial spin state is often not clear, which prevents

further interpretation of the coupling between spin and other ordering

parameters at oxide heterointerfaces.

A good candidate for the magnetic constituent of such interface is

doped manganite, which received detailed understanding on carrier filling and

orbital effects [153, 154], To realize electronic and structural reconstructions

of doped manganite, electrostatic and strain effects are primary methods,

which modulate the competition between different interactions [155-160], In

this letter, the spin state at the PZT/LSMO interface is directly probed by the

interface-specific MSHG technique, and the control of exchange coupling

between interfacial spins is realized through their coupling with the charge

degree of freedom, which provides previously unknown critical insight into the

magnetic order at this multiferroic heterointerface.

118

6.2 Samples and experiments

The PZT/LSMO heterostructures used in this study are grown by

pulsed laser deposition on (100) LaAI03 (LAO) substrates. The LSMO layer

is grown at 600 °C under an oxygen pressure of 80 mTorr, using a laser

energy density of 1.8 J/cm2 and repetition rate of 10 Hz, followed by

annealing at 700 °C for 30 min in oxygen. The PZT layer is then deposited on

the LSMO layer at 600 °C at the same conditions. After deposition the

heterostructure is annealed at 700 °C for 30 min in oxygen at a pressure of

300 Torr. Finally, the films are cooled down to room temperature at a slow

rate. The film thickness of PZT and LSMO layer is around 550 nm and 100

nm respectively, determined from an X-P-200 profilometer. FE PZT layer is

used as a constituent of the heterostructure to adjust the carrier injection at

the interface layer, considering its high Curie temperature Tc [161], high

remnant polarization and low coercivity field £ c [162]. Such thick PZT layer is

used to suppress leakage current. The ferromagnetic LSMO layer is known to

exhibit high Curie temperature of about 370 K, colossal magneto-resistance

properties and half metallic behavior [155,163], and typically is used as a

spin-polarized detector [164], Indium tin oxide (ITO) top electrode is deposited

on top of the PZT layer to enable the application of an electrical field.

To investigate the interfacial magnetic state of the PZT/LSMO

heterostructures, we carry out MSHG measurements with fundamental laser

beam at 800-nm wavelength and detection of MSHG signal at 400-nm

wavelength, as depicted in Fig 6.1(a). The MSHG signal can be generated in

119

magnetic materials in which both space-inversion and time-reversal symmetry

are simultaneously broken [26-28, 165-167]. The broken time-reversal

symmetry results from the presence of magnetization, while the broken

space-inversion symmetry occurs only at the surface or interface of

centrosymmetric materials like LSMO. Considering these requirements, the

first unit cell layer of LSMO at PZT/LSMO boundary can generate MSHG

signal. However, the Nth (N>2) unit cell layer from such boundary will not

generate MSHG signal, as it is surrounded by the (N-1)th and (N+1)th unit cell

layers which preserve space inversion symmetry (Fig. 6.1(b)).

For comparison, MOKE measurements are carried out to probe the

magnetic state of LSMO bulk layer. MOKE technique is carried out with

similar experimental configuration as MSHG (Fig. 6.1(a)), except that the Kerr

signal (800 nm) is detected with a high-speed silicon detector after a P-

polarized analyzer. The optical absorption coefficient in LSMO for SHG light

(A = 800nm) is about 1 * 105cm-1 . Considering the exponential decay of light

intensity, the magnetic contribution of the first three unit cell layers to the

MOKE signal which probes the magnetization of the entire LSMO film is less

than 2%, which is close to the noise level of the MOKE hysteresis loop.

Hence, the MOKE technique is not sensitive at all to the change of interface

magnetization.The external magnetic field H is applied along the [100] axis of

LSMO layer for both MOKE and MSHG.

Figure 6.2(a) shows the X-Ray Diffraction results of PZT/LSMO

heterostructure measured with a Siemens D500 x-ray diffractometer (Cu Ka

120

radiation) in a 0-20 scan. Only the (100) diffraction peaks of LSMO and PZT

are observed, along with those of the single crystal LaAI03 substrate. This

suggests that the individual ferromagnetic and ferroelectric layers are grown

epitaxially block-by-block on the substrate and form smooth films without any

secondary parasitic phases. The LSMO layer is under slight tensile strain with

lattice parameter a=3.885 A0, due to the oxygen pressure during growth [168,

169]. Figures 6.2(b)-6.2(d) show RHEED patterns of LAO substrate and

LAO/LSMO structures before and after the deposition of PZT layer on top.

Very intense specular spots, Kikuchi lines and Laue zones are observed,

revealing that the LAO substrates used in this work are high-quality single

crystals and perfect in-plane epitaxial growth of the LAO, LSMO and PZT

layer with smooth surface.

121

(a) MSHG MOKE

PZT

Interface

LSMO

• Ti Zr

PZT

Interface

LSMO

Figure 6.1 Schematic of MSHG and MOKE measurements (a), and depiction of

PZT/LSMO interface on atomic scale (b). MSHG probing plane is marked with purple

color.

122

oo

CM

20 30 40 50Braggs Angie (2e)

(b) (c) (d)

LAO RHEED pattern ISMO/LAO RHEED pattern PZT/LSMO/LAO RHEED pattern

Figure 6.2 XRD results of PZT/LSMO heterostructure (a), RHEED patterns of LAO

substrate surface (b), LSMO surface before deposition of PZT layer on top (c) and

PZT surface of PZT/LSMO/LAO heterostructure (d).

123

6.3 Charge control of antiferromagnetism at PbZr0 .s2Ti 0 .4 8 O3/

Lao.6 7Sr0.33Mn0 3 interface

Figures 6.3(a) and 6.3(b) display the interfacial and bulk magnetization

as a function of applied gate voltage Ug probed with MSHG and MOKE at 78K,

respectively. The key finding is that the magnetization of the interfacial layer

is modulated by Ug, while the bulk magnetization is not. The magnetic

contrast for MSHG and MOKE is defined as

„ _ /(+M )-/(-M ) /cA ~ r(+M)+n - m ( 6 1 )

where l(±M) is the signal intensity received by the detector when the

macroscopic magnetization M is aligned with external magnetic field in either

positive or negative [100] direction (Fig. 6.4(a)). The magnetic contrasts

obtained by both MSHG and MOKE is linear with M in the probed volume, if

magnetic component is much smaller than non-magnetic component in signal

intensity as in our case [26], The magnetic contrast A of MSHG measurement

as a function of Ug is displayed in Fig. 6.4(b), where positive Ug refers to

positive voltage applied to the ITO top electrode. These results indicate that

the modulation of gate voltage on PZT/LSMO heterostructure, which is not

strong enough to change the magnetization of the bulk LSMO layer, can

significantly tune the magnetic state at the interface. This may provide a clue

for fabricating low energy consumption multifunctional devices based on

similar material systems.

124

interface10 10

0 P o

7.5 V -7.5 V

-10 V -10 V- 10-10

-1000 1000 -1000 1000Magnetic field (Oe) Magnetic field (Oe)

Figure 6.3 Interface (a) and Bulk (b) magnetic hysteresis loops from PZT/LSMO

heterostructure for different external gate voltages Ug measured with MSHG and

MOKE techniques at 78 K, respectively.

125

MOKE

sig

nal (

a.u)

0.36

3.0.32

0.28

0.1

<

0.0

Figure 6.4 MSHG hysteresis loop with Ug = +10V indicating l(-M) and l(+M) used to

determine the magnetic contrasts from Eq. (6.1) (a), and MSHG magnetic contrast

A as a function of Ug (b). Increasing (decreasing) gate voltages, Ug up (Ug down), are

labeled in red (black). The direction of positive Ug is defined in the inset.

Ug down Ug up

Ug (V)

I (+M)

l(-M)

1000-1000 0H (Oe)

126

(a)

Ug down Ug up

-60

-15 0 15Ug(V)

(b)

Ug down Ug up0.1

0.0

60-60 0

P (jaC/cm2)Figure 6.5 Ferroelectric polarization (P) in PZT as a function of Ug (a). The direction

of positive P is defined as pointing towards LSMO thin film. M SHG magnetic contrast

A as a function of P of PZT (b). The dashed line in Fig. 6.5(b) is the least square fit

for the data points.

127

(a)

OOo

CO

4

-10

(b)

0ug(V)

10

Ug down Ug up

- + — --

0.0

3 4S (1/1000)

Figure 6.6 Strain state (S) of LSMO interface as a function of gate voltage (a). MSHG

magnetic contrast A as a function of strain (5) of PZT (b).

128

To elucidate the change of interface magnetization with gate voltage,

FE polarization (P) measurements are carried out, as shown in Fig. 6.5(a).

The FE polarization saturates with gate voltage above 10V (electric field of

approximately 180 kV/cm) or below -10V, indicating that all FE domains are

aligned. The switching transition around Ug = ±2V is not sharp, resulting from

the variance in switching threshold of different FE domains [170]. The hole

density at the PZT/LSMO interface can be electrostatically modulated, i.e.

hole depletion/accumulation occurs when P points towards/away from the

LSMO layer [149-151]. From the FE polarization measurement, we derive the

change in magnetic contrast A as a function of P, as shown in Fig. 6.5(b),

using the dependence of A on gate voltage Ug (Fig. 6.4(b)). The relationship

between A and P exhibits a linear correlation, independently of the history of

polarizing PZT layer, as the data with decreasing gate voltage (black)

overlaps with the increasing one (red).

The lattice constant of PZT layer can also be simultaneously

modulated by gate voltage application, due to the piezoelectricity of PZT. To

estimate the influence of gate voltage (Polarization) on interface strain state,

we used the lattice distortion of PZT layer from Ref. [171]. In this estimation,

we assumed that the modulated x-y plane strain of PZT by gate voltage is

transferred onto LSMO lattice at interface. This assumption might lead to a

larger estimation of LSMO lattice distortion, but is still similar with realistic

case. The strain S is defined as the percentage variation in the in-plane

lattice parameter of LSMO layer relative to the bulk LSMO value, and

129

displayed in Fig. 6.6(a) as a function of gate voltage. The offset at zero

voltage (about 0.3%) is based on the lattice constant derived from XRD

measurement. The strain S at interface is weakly modulated by gate voltage

Ug, and its dependence on gate voltage Ug exhibits a butterfly shape.

Therefore, the magnetic contrast A as a function of strain S is depicted in Fig.

6.6(b), which is derived from the relationship between strain and gate voltage

in PZT thin films studied in Ref. [171] and using the dependence of A on gate

voltage Ug (Fig. 6.4(b)). There is no clear correlation between interface

magnetization and strain.

Thus, the observed change of interface magnetization with gate

voltage results from charge injection. The charge mediated mechanism is

depicted by a schematic model of screening charge accumulation at the

PZT/LSMO interface (Fig. 6.7), responding to opposite FE polarization

orientations in the PZT layer. The polarization follows the relative

displacement along the z direction of the metal cations (Pb, Zr and Ti) with

respect to their neighboring oxygen anions in the same (001) plane [151].

When the FE polarization of PZT is pointing towards the LSMO layer,

electrons will accumulate at the LSMO surface to compensate for the charge

on the PZT side. On the other hand, when the FE polarization of PZT is

pointing away from the LSMO layer, holes accumulate at the surface of

LSMO. To quantitatively determine the hole doping level p modulated by gate

voltage at the first unit cell layer of interface LSMO, we assume that the

Polarization (P) of PZT initiates the charge at geometric boundary between

130

PZT and LSMO (z=0) as P 5 (z ) , and the density of screening charges

accumulating at the LSMO side of interface exponentially decays with the

distance away from PZT/LSMO geometric boundary, as depicted in Fig. 3.

Thus the charge distribution can be written as P8(z) e~z/x a t z > 0 ,A.

which is at the LSMO side of interface. The Thomas-Fermi screening length

X = yje/NEfe2 = 4.14 >4°, where s = yj32e0 is the dielectric constant, NEf =

0.6(ev~1u.c.~1) is the Fermi level density of states [172]. The modulation of

hole doping level per Mn ion at single-unit-cell interface LSMO layer is

approximately calculated from the integration of charge density from z=0+ to

z= lattice constant of LSMO along [001] direction, and the degree of which is

calculated to be ±0.231 electrons per Mn ion in the first unit cell layer at the

interface corresponding to ± 10V gate voltage.

PZT(z>0)

e l i p r - @y LSMO/ (? < 0) 1

(z = 0)

Figure 6.7 schematic model of ions displacement and screen charge accumulation

under opposite polarization state.

The loss of interface magnetic contrast A by hole injection (negative Ug)

can be due to either a modulation of saturation magnetization, or changes in

the coupling mechanism between Mn ions, or changes in the

magnetocrystalline surface anisotropy, or a combination thereof. The spin

alignment of Mn 3d electrons still obeys Hund’s rule, and the magnetic

moment per Mn site cannot be smaller than 3 Bohr magnetons, as the charge

injection by PZT can only modulate the eg electron concentration [149, 172],

Thus, the depletion of electrons with coherent spins by hole injection can

cause a reduction of magnetization, but the degree of which is limited and

thus not responsible for the significant change of interface magnetic contrast

by one order of magnitude.

Another possibility is that the surface magnetic anisotropy changes

with increasing hole doping and the magnetization starts to point out of plane.

In the probe with optical polarization configuration S in S out, the magnetic

tensor X y y y ( M x ) is only sensitive to the Magnetization changes in the

longitudinal direction, and can not exclude such possibility. In addition, we

carried out MSHG measurement with P in S out optical polarization

combination, where magnetic tensor contains both XyzziMx) . Xyxx{Mx)

sensitive to longitudinal magnetic moment and Xyzx(M z) sensitive to polar

magnetic moment. The results under opposite polarization states of PZT are

displayed in Fig. 6.8. Under -10V gate voltage, there is no magnetic contrast,

hinting not only M x but also M z become zero. There is still a possibility that

surface magnetic anisotropy change as hole doping is modulated. But in our

132

case, we pull the interface macroscopic magnetization aligned with strong

enough magnetic field. This speculation is also supported by the fact that the

magnetization in magnetic hysteresis loop is saturated above 300 Oe. Thus,

we conclude that the interface magnetization is reduced to zero when

negative gate voltage is applied.

t 136 ’</> tz0)

■M

i 1.32cg>'</)<D S 128

-600 0 600 H (Oe)

Figure 6 .8 P in S out M SHG results under different polarization states.

Therefore, the only reason is a change of exchange coupling at the

PZT/LSMO interface by hole injection as pointed out by Vaz et al. [149]. The

diminishing magnetic contrast in MSHG measurement with increasing hole

doping suggests the appearance of AFM spin alignment in the first unit cell

layer at the PZT/LSMO interface.

—- +10 V U ■ "1"1 * ■— -10V 1 k

133

Next, we discuss the evolution of interface spin alignment in terms of

3d orbital occupancy of Mn ions. The origin of the AFM phase at the interface

is not completely understood. A recent discovery shows that d3z2 _ r 2 orbital

has lower energy than d x2 _y2 orbital, and thus is more occupied with electrons

at tensile strained LSMO surface or interface as a result of broken space

inversion symmetry [173-176], in contrast to tensile strained LSMO bulk.

According to Ref. [173], the d x2 _ y2 orbitals are unoccupied at the interface,

which results in anisotropy of exchange coupling [148]. Double Exchange

induced by hopping of eg electrons is stronger in the z direction (surface

normal), while super exchange induced by local t2 9 electrons is stronger in the

x-y plane (surface plane). It is demonstrated by C. Aruta et al. [175] that this

interfacial d 3z2 _ r 2 orbital occupation favors the C-type AFM spin ordering, and

the easy axis of the AFM phase is oriented along z direction, irrespective of

the strain conditions. Moreover, A. Tebano et al. [173] provided microscopic

evidence that this preference in orbital occupancy at the surface of LSMO

exists, independent of the chemical nature of the substrate and the presence

of capping layer. This is consistent with our observation that spin coupling on

x-y plane is tuned into AFM type at the PZT/LSMO interface, as the d x2 _y2

orbital becomes depleted by hole injection. Furthermore, as electrons

accumulate at the LSMO interface layer by tuning the FE polarization of the

PZT layer, in-plane orbital occupation of eg electrons is also expected,

corresponding to the appearance of FM coupling on the x-y plane. Further

injection of electrons into interface layer might lead to expansion of A-type

134

AFM regions at the expenses of the FM double exchange regions, which

resembles the AFM coupling of LaMn03[169]. Further studies are required;

particularly in view of the fact that spin transport properties may depend upon

the phase transition by charge injection or oxygen vacancy concentration at

the interface.

The transition from AFM to FM spin alignment at the PZT/LSMO

interface with increasing P is gradual (Fig. 6.5(b)), indicating that multiple

domains of both phases coexist as the ferroelectric polarization is switched.

Multiple domains with different polarizations exist near the gate voltage where

FE switching occurs in PZT thin film, as reported in Ref. [170]. Some FE

domains may be switched to the opposite polarization, which leads to a

distribution of local FM and AFM domains in the LSMO interface layer. When

P is saturated and points towards (away from) the LSMO film, the interface

spin alignment is FM (AFM) coupled. Thus, the redistribution of local

magnetic domains at the PZT/LSMO interface leads to a gradual change of

magnetic contrast in MSHG measurement.

135

Chapter 7

Conclusions

Fundamental understandings of magnetic properties and spin

dynamics help facilitate the design and fabrication of new magnetic devices

with desirable performance. In this dissertation, intrinsic Gilbert damping a0

and perpendicular magnetic anisotropy (PMA) in L1oFePd(i.X)Ptx ternary alloy

films can be continuously tuned by spin-orbit coupling strength f with

appropriate Pt/Pd concentration. In particular, aro and PMA are found to be

proportional to <f2 from TRMOKE studies for the first time, which is consistent

with ab-initio density functional calculations. The tunability of PMA is

enhanced by increasing chemical order in FePdi.xPtx alloys and by lowering

the temperature. Moreover, control of the anti-site disorder c with proper

growth temperature results in significant variation of aro in FePt thin films. As c

increases from 3% to 16%, aro increases by more than a factor of three from

0.06 to 0.19. The variation of c mainly affects the electron scattering rate 1/re,

while other leading parameters remain unchanged. A linear correlation

between croand 1/re is experimentally observed, consistent with the spin-orbit

coupling torque correlation model for interband transitions. The ar0 remains

unchanged with temperature, indicating that the intraband contribution to aro is

inundated by impurity scattering. Furthermore, as anti-site occupation

decreases, the perpendicular magnetic anisotropy increases and the Landau

136

g factor exhibits a negative shift due to an increase in orbital momentum

anisotropy.

The efficiency of spin precession excitation by laser pulses is

significantly improved through inserting an AFM CoO layer into Fe/MgO

heterostructure which establishes the uniaxial magnetic anisotropy Ku along

Fe [100] direction. The modulation of Ku by laser pump pulses generates a

fast exchange torque to the FM Fe magnetization. The transient torque is

enhanced at temperatures where Ku varies significantly and with external

magnetic fields competitive to the magnetic anisotropy fields, which results in

spin precession amplitude peaks. The excitation approach by modulation of

FM-AFM exchange interaction is much more efficient with 400-nm-

wavelength pulses via photo-excited charge transfer processes in AFM CoO

layer than modulating the FM order of Fe with 800-nm-wavelength pulses.

The recovery time of the exchange torque is around 40 ps, much faster than

the cooling time from demagnetization. Our results will help promote the

development of low energy consumption magnetic device concepts for fast

spin manipulation at room temperature.

This dissertation also presents photo-induced magnetization

precession in half-metallic Lao6 7 Cao3 3Mn0 3 thin films grown on NdGa0 3

substrates with TRMOKE technique. For the first time, pronounced spin

precessions are observed in a geometry with negligible canting of the

magnetization with a wide range of magnetic fields (0-2.5 T) applied along

the in-plane easy axis. The precession amplitude monotonically decreases

137

with increasing field, in contrast to the precession induced by the transient

anisotropy field in the state of canted magnetization. These observations

indicate that the coherent spin rotation may be triggered by a transient

exchange interaction and not by demagnetization and/or anisotropy field

modulation. The picosecond modulation of the exchange interaction results

from the emergence of AFM interaction caused by charge transfer and

alteration of kinetic energy of eg electrons in LCMO under optical excitation.

Finally, the interfacial spin state of PZT/LSMO heterostructure and its

dependence on FE polarization is investigated with magnetic second-

harmonic generation at 78K. The magnetization in the first unit cell layer at

the interface can be widely tuned over an order of magnitude by varying the

applied electric field between ± 1 8 0 kV/cm. The cross-coupling between

ferroelectric and ferromagnetic behavior at the heterointerface is mainly

electronic, due to charge injection. The ferroelectric polarization of PZT can

tune the spin alignment of Mn ions at the heterointerface from FM to AFM

exchange coupled. In contrast, the bulk magnetization probed with MOKE

remains unchanged. Multiple domains of FM and AFM phase coexist as the

ferroelectric polarization of PZT is switched.

138

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VITA

Xin Ma was born in Fuzhou, Jiangxi, China on May 18, 1988. Xin Ma

received his B.S. degree majored in Condensed Matter Physics at the

University of Science and Technology of China in 2009.

In August 2009, Xin Ma entered the College of William and Mary as a

research assistant in the Department of Applied Science. As a PhD candidate,

Xin Ma devoted to investigating magnetic properties and materials, including

spin dynamics and interface magnetization, with pulsed lasers and ultrafast

techniques.

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