W&M ScholarWorks
Dissertations, Theses, and Masters Projects Theses, Dissertations, & Master Projects
2015
Optical characterization of ferromagnetic and multiferroic thin- film heterostructures
Xin Ma College of William & Mary - Arts & Sciences
Follow this and additional works at: https://scholarworks.wm.edu/etd
Part of the Artificial Intelligence and Robotics Commons
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
This Dissertation is brought to you for free and open access by the Theses, Dissertations, & Master Projects at W&M ScholarWorks. It has been accepted for inclusion in Dissertations, Theses, and Masters Projects by an authorized administrator of W&M ScholarWorks. For more information, please contact [email protected]. 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 Technology of 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 ar 0 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 disorderc 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 7 Cao.3 3 Mn0 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 2 Tio.4 8 O 3/Lao.6 7 Sro.3 3 MnO3 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 harmonic generation 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 in FePtxPd(i-x) thin films 43
3.4 Magnetic anisotropy and damping in FePtthin films 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 photo excited 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 7 Cao.3 3 Mn0 3 103
Chapter 6 . Magneto-electric coupling at PbZr 0 5 2 Ti 0 .4 8 O3 / Lao 6 7 Sr0.3 3 MnO3 interface
6.1 Introduction 117
6.2 Samples and experiments 119
6.3 Charge control of antiferromagnetism at PbZr0 52Ti 0 4 8 O 3/ 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 incident angle, 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 TRMOKE measurements for different doping level x at 20K and 200K. LIST OF FIGURES
1.1 Schematics of magnetic recording arrays with traditional writing mechanism. 2
1.2 Schematic of precessing magnetization. 3
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
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). 15
2.2 Geometry of TRMOKE measurements. 17
2.3 Stages of the excitation process: (a) M aligns along the equilibrium direction H eff \ (b) new equilibrium direction H'eff when pump pulse “heat up” the sample; (c) M precesses around H'eff (d) heat diffused and M processes around H eff. 19
2.4 TRMOKE result from L10 FePt thin film with growth temperature 550 °C at magnetic field 3.5T. 20
2.5 Schematic of TRMOKE measurement geometry for L1 o FePt thin films. 21
2.6 Coordinate system used for calculating the magnetization precession frequency in the L10 FePt thin films. 22
2.7 Schematic of TRMOKE experimental set-up. 25
2.8 Schematic of MSHG and MOKE measurements on PZT/LSMO thin film. 26
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). 31
2.10 The schematic set-up of MSHG measurement. 32
3.1 Schematics of atomic structures for the two series of FePt thin films. 37
vi 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. 39
3.3 Anti-site disorder percentage c in Z_10 ordered FePt thin films as a function of growth temperature Tg. 40
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
3.5 TRMOKE data with different doping level of FePtxPd(l-x) under magnetic field H=5T. The solid lines are fitted curves. 45
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. 46
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. 47
3.8 PMA (a) and ao (b) as a function of doping level x of FePtxPd(i-x) 48
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). 49
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. 53
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. 56
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. 57
vii 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
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. 61
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. 6 5
4.1 Longitudinal hysteresis loops for magnetic field along uniaxial easy axis [100] and hard axis [010] at 82 K and 330 K. 70
4.2 Temperature dependence of coercivity ( Hc) and exchange biasing field ( He) obtained from easy axis loops. 71
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. 73
4.4 Spin precession frequency f as a function of magnetic field H. 74
4.5 Temperature dependent TRMOKE results with magnetic field H= 3 kOe. 76
4.6 Temperature dependent TRMOKE results with magnetic field H= 1 kOe. 77
4.7 Spin precession frequency fa s a function of temperature. 78
4.8 (p as a function of temperature (green=3 kOe, blue=1 kOe). 79
4.9 9 as a function of temperature (green=3 kOe, blue=1 kOe). 80
4.10 Uniaxial magnetic anisotropy field KJMs as a function of temperature. 81
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. 85
viii 4.12 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. 86
4.13 Measured and simulated spin precession amplitudeA as a function of T. 88
4.14 The spin precession amplitude A is plotted as a function of H at different T, where solid lines represent simulated results. 89
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
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. 95
4.17 Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with 77 = 2 ps. 96
4.18 Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with 7> = 15 ps. 97
4.19 Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with rr = 1 0 0 ps. 98
5.1 Transient reflectivity A R (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). 104
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
5.3 Three-temperature model in ferromagnetic metals and half-metals. 106
ix 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
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. 111
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
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
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 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
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. 126
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. 127
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). 128
X 6.7 Schematic model of ions displacement and screen charge accumulation under opposite polarization state. 131
6.8 P in S out MSHG results under different polarization states. 133
xi 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].
*Ring* writing element
Longitudal Recording (standard)
] Recording layer
'Monopole’ writing element
Perpendicular Recording
L a Recording Layer | Additional Layer
Figure 1.1 Schematics of magnetic recording arrays with traditional writing mechanism [16].
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) dt 1 M dt ' ’ 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 eff ■ 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
AFM
(ii)
-> H (iii)
(Hi) (iv) (iv)
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.
8 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 7 Cao3 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 7 Ca0.3 3 Mn0 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 2 Tio 4 8 O3
/La0 6 7 Sro.3 3 Mn0 3 and corresponding experimental consideration for probing
interface magnetization.
11 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 -iQz 1 iQx (2 .1) iQy iQx 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
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 ofnL 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:
(2.2)
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
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
* ie w + * * i. ' ' j-i, * wt 1 ,A' ■ . * a1-na2 •. jf(naa+az)(na2+ a i) a t+ n a 2
Longitudinal n a 1- a 2 Qf,na1tan02 «1-n a 2 n a x+ a 2 i(na1+a2)(na2+a1) ax+na2
-WW... l^C ^/o^ai—thQ^itanBz ■ * ax-na2 ax+na2 i ;t »£} pry*^ ' ^ T *?~~ ' . /.'.w <• I / “ • I; X ; i l
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.
_ Pump Probe TRMOKE
Substrate
Figure 2.2 Geometry of TRMOKE 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).
to < t< to + 1 0 ps t> to + 1 0 ps
Figure 2.3 Stages of the excitation process: (a) M aligns along the equilibrium direction H eff \ (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 40 Delay time (ps)
Figure 2.4 TRMOKE 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 byHeff, 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 TRMOKE 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 = 4tc Ms . The equilibrium direction for M can be derived by taking the derivative of F over 0:
n _ sin(20) Ha-H d 2*sin((p-6) ' ' ’
i L Z ’ //[001] z ///II
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:
-» -* X y z d m x •* , d m y d m z —~x + ——y + — mx my Ms + mz (2.6) d t d t d t y Hx Hy Hz where Hx,Hy,Hz are the x, y, z components of Hefr, and can be derived by:
dF d M y
Hy = H s\n(0 - qi) (2.7)
dF + He os(0 - Therefore, the torque equations become: 1 drriy = — [H cos(0 — ^ = [H cos (0 - For the uniform precession, the my and mz can be written as: m y = m yQel2TCft+ i2nft+ The precession frequency f of spin wave can be derived by substituting Eq. (2.9) into Eq. (2.8): 2tt/ = yV[W cos(0 - (2.10) 23 In terms of Gilbert damping parametera, 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 - (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/cm 2 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. Analyzer (S) Polarizer (P) Electromagnet Figure 2.7 Schematic of TRMOKE 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 2 Tb 4 8 O 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/LSMO thin film 26 2.3.1 Second harmonic generation (SHG) SHG signal arises from the nonlinear polarizationP(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>) = 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> Xzyy 3nd Xzzz< which are 28 even terms, x Xxy> Xxyx> Xyxx> Xyyy> Xyzz> Xzyz Xzzy 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. Input Output Polarization Polarization yzz, yxx P s Longitudinal zzz, xzx PP xxz, zxx Jlf//x zyy s P yyy s s P s Transverse xxz, xzx zxz, zzx P p zxx, zzz XXX, xzz H f//y zyy xyy s p s s yxz, yzx p s Polar xxz, xzx p p zxx, zzz M//z zyy s p s 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\ Xeff ^li,j.k&ijkXijk (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 1000 Magnetic Field (Oe) Figure 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). 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-ps filtei Prism Lens Lens Photodiode Polarizer Analyzes 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) Varied chemical order modulating spin-orbital coupling 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' measmeas\ /o ('001) (3.1) '^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/Ioo2)caic *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=0.5 x=0.25 & to c Q xa : fct (001) fct (002) 20 30 40 50 60 20 (deg) (b) fct (001 ) fct (002) Tg=720 °C f j u S | Tg=680 °C (/) c P A 0 Tg=620 °C c I I A o a Tg=580 °C a: X MgO J L 1 1 i ...... 20 40 60 20 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 1 6 ••••••• • • • • • • * FeJ ••••••• • ••••• • Pti 12 ••••••• 8 2 2 nm FePt o 17 nm FePt •••••••• !•••••• f t • -'••••••• ••••••• • Pt J ••••••• 480 540 600 660 720 Tg(°C) 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) x=0 — 0 . =o -- 0 =90 H (T) 0 0 0 0 6 0.0004 0.0002 0.0000 r - 0.0002 -0.0004 -0 0 0 0 6 -3-2-10123 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 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 = > 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 ofg 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=0.5 zj x=0.25 CD =0.15 10 20 30 40 Delay t (ps) Figure 3.5 TRMOKE 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 T CD H= 2 T 10 20 30 40 Delay t (ps) * X= 0 5 H= 5 T j / V \ A a a a a a a / w w i/VV v v ^ ^ zs CD o' H= 2 T 20 40 60 80 Delay t (ps) Figure 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. 46 0.3 0.2 N X I- X=1 x=0.9 x=0.75 x=0.5 0.0 x=0 0 2 4 6 H (T) 0.6 • x=0.9 ▲ x=0.75 0.3 ▼ x=0.5 ♦ x=0 0.04 0.00 0 2 4 6 H(T) Figure 3.7 The dependence of precession frequency f (a) and lifetime x (b) on external magnetic field Hfor typical samples. The solid lines refer to fitted results. 47 n ^ 4 CO E ,0 ^5) - 0 r^- O 2 — / T— i zs • ■ ■ ■ 0 I i . i 0.0 0.5 1.0 X Mb) 0.10 0.05 0.00 0.0 0.5 1.0 Figure 3.8 PMA (a) and a0 (b) as a function of doping level x of FePtxPd(1.x). 48 0.6 (a) >d) 0.4 ■ 0.2 I 15 -(b) > 10 5 5 - (c) CD3.0 ZL Fe 1.5 U) Pt/Pd F 0.0 ■ ■ 1.5 '(d) <0 1.0 O 0.5 __L_ 0.0 0.5 1.0 x 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 ar0at 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 10 ■ measured 8 • calculated E .o 6 13> ♦ chem ordered Q) 4 O 13 2 0 0.0 0.3 0.6 SOC ^ (ev) measured 0.10 s° 0.05 0.00 0.0 0.3 0.6 SOC $ (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 H k(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 H k (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 fa s 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 (a) c= 4.0% c= 8.6% TO,13 * c= 10.5% c= 13.8% 10 20 30 t(ps) c= 4.0% H= 6 T H= 5 T =3 d H = 4 T H= 3 T 20 30 t(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 differentH (b). The solid lines refer to fitted results. 56 0.35 c= 4.0% c= 8.6% c— 10.5% 0.30 c= 13.8% N I I- - 0.25 0.20 24 6 H (T) 0.20 ▼ c= 13.8% a c= 10.5% • c= 8.6% ■ c= 4.0% 0.15 4) 0.10 0.05 2 4 6 H (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 (a) E o E> a> h- o ■ 22nm FePt =3 - 17nm FePt 0 6 12 18 Anti-site percentage c 2.2 o o £ 2.0 o> 0 6 12 18 Anti-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 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., 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 Using the fitted values of r, we determinea 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 (a) 22nm FePt 17nm FePt o> ■£= 0.14 Q. E (0 ■O ■C a> S 007 0 6 12 18 Anti-site percentage c 0.09 22nm FePt 17nm FePt 0.06 o> 0.03 0 6 12 18 Anti-site percentage c Figure 3.14 a0 (a), and or 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/M02A« (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 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 (a) 0.3 0.3 Frequency Decay rate N 0.2 0.2 £ 0.1 - • 0.1 0 10 0 200 T (K) 0.30.3 ■ Frequency 0.2 • Decay rate N X I— 3 4 5 6 H (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. CO m mmm 8 2 K c H / / [ 1 0 0 ] ■ -Q 3 3 0 K co H / / [ 1 0 0 ] co c o> 0 ) CD * - 1.0 - 0.5 0.0 0.5 1.0 Field(KOe) Figure 4.1 Longitudinal hysteresis loops for magnetic field along uniaxial easy axis [100] and hard axis [010] at 82 K and 330 K. 70 600 T ------1------1------,------1- He ^ 4 0 0 He 0 £200 o 6 o 1 o 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 = yJ[Hcos(S- f) + Ha][Hcos( Ha =Hdeff- 2KuS'm2flM s + K,(2- sin22 H? = 2K,cos40M s + 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 72 •— 2112 ■ 1848 V ■ 1320 — 658 * * ► 461 ■ 330 264 05 • 198 — 132 >— 66 — minus132 - 0.1 0.0 0.1 0.2 0.3 0.4 0.5 Delay time t (ns) Figure 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. 73 20 N X o -1000 0 1000 2000 3000 H (O e) Figure 4.4 spin precession frequency fas 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 2co s26 + 2nMs) — ^ (sinBcoscp + sinQsin (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 = y J W J h (4.4) 75 300K 290K 280K j * ‘ . * •* *''** .*—* »*S4^ 270K 3 7 \ V a ^ - ^ w 260K nj, r \ A U * U * ~ ^ 250K lc s .-. •. ; c ^ * ♦/« y ***^ 240K g» * * 4 « 4 4 ^ ^ _ 230K to 210K a) 190K * 150K c 100K a) -t._/\ yv y \ v \ - / ,%- ' 50K c 10K CO 0.0 0.1 0.2 0.3 0.4 0.5 t(n s ) Figure 4.5 Temperature dependent TRMOKE results with magnetic fieldH= 3 kOe. 76 Figure 4.6 Temperature dependent TRM O K E results with magnetic field field magnetic with results E K O TRM dependent Temperature 4.6 Figure Transient Kerr signal (a.u.) • . / ■. • • 77 ...... 50K 250K 260K 270K 280K 290K 210K 230K 240K 10K 100K 300K 150K 190K H= kOe. 1 35 « 30 _ 25 N X o H = 1 kOe H = 3 kOe 20 15 0 50 100 150 200 250 300 T( K) Figure 4.7 spin precession frequency f as a function of temperature T. 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(Mcos#)2 4 (4.5) - (sin 0 cos Along the equilibrium direction of magnetization, F reaches the dF dF minimum, 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 - 100 150 200 250 300 Temperature (K) Figure 4.8 q> as a function of temperature (green=3 kOe, blue=1 kOe). 79 Figure 4.9 4.9 Figure theta (degree) 84.5 85.5 86.5 87.5 88.5 G a ucino eprtr gen3ke bu= kOe). blue=1 kOe, (green=3 temperature of function a as 100 eprtr (K) Temperature 80 150 200 5 300 250 2.0 0 O Ku/Ms 0.5 0.0 0 100 200 300 T (K) Figure 4.10 Uniaxial magnetic anisotropy field KJMS 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 The amplitude a qualitatively describes how far the FM system is driven away from the ground state, and the amplitude ratioA/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 temperatureTe 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* ®@®@ FM Fe - @ ® @ ® @@@@ FM/AFM XX exchange i Co 3d — AFM CoO — hv~ 3.leV •••• 3 > 0 2 p 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 - Fe/MgO 3.1 mJ/cm2 • Fe/CoO 3.1 mJ/cm2 a Fe/CoO 0.16 mJ/cm2 =3 aj • \ j l ,A=800nm ro c a 'to ^ 'A A A A a a k. fc A=400 nm * 0 200 400 Delay t(ps) Figure 4.12 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. 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 20 H = 2 kOe < 10 0 1 0 0 200 300 T(K) Figure 4.13 Measured and simulated spin precession amplitudeA as a function of T. 88 T= 240K T= 300K T= 78K 0 0 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 O e*Ms), which agrees quite well with the measured relationship betweenA 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- AKu(t+At))- 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. 0.02 step size = 0.8 ps xr = 40 ps £ 0.01 s 0.00 o> - 0.01 AAT, - 0.02 0 40 80 Delay 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 step size = 4 ps xr = 40 ps 0.01 3 05* CO 0.00 "ro §N c E g> co - 0.01 a> - 0.02 0 200 400 Delay 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. 95 step size = 0.8 ps 0.001 Tr = 2 ps ^ ni 0.000 Kerr signal (a.u.) signal Kerr - 0.001 0 40 Delay t (ps) Figure 4.17 Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with rr = 2 ps. 96 0.02 step size = 0.8 ps Tr = 15 ps / y * 0.01 * 0 . 0 0 Kerr signal (a.u.) signal Kerr - 0.01 0 40 80 Delay t (ps) Figure 4.18 Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with rr = 15 ps. 97 _ - step size = 0.8 ps 0.02 — xr = 100 ps 0.01 CO ^ 0.00 £ - 0.01 Kerr signal (a.u.) signal Kerr - 0.02 0 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. 98 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 7 Ca0 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 7 Ca0.3 3 MnO3 (LCMO) thin films on NdG a03 (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 H I I Easy Axis (c) 3 m 5 0.05 T 0 50 100 150 0.3 T Time (ps) H x Plane (b) 3 ■ <3 ■ V 0 100 2 0 0 50 100 150 200 Time (ps) Time (ps) 3 % H = 0 T 0 200 400 600 0 Time (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 hard axis (a) easy axis (b) 560 Oe « 560 Oe OOe 0 200 400 0 200 400 Time (ps) Time (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. 1 (a) 0.9 -O . cSrEPn . LCMO 2 ■D a hard axis easy axis ■ □ □ 0.6 ■%9 0 □ o ,S 60 nm 0.3 X i 1 i 1 1.2 o □ « □ n D □ o 3 • □ 0.8 o ^ 1 0 0 nm o 0.4 1.2 □ o □ ■ ■ ■ □ 0.8 o . □ o g 100 nm* □ 0.4 OO ...1.... ■... .. A-...... L J— ...... 1..... -..... 1... 1 2 0 Magnetic Field (T) (b> o 0 o 0 □ 1 £ 1 • 1 □ 6 □ □ 1 cs T> r □ o Fe — □ ■' 4 -o O. hard axis . q If f to easy axis E □ 2 < I * ....A - 1 1 . 1 i 1 0 400 800 0 400 800 Magnetic 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 50 o> Z 40 30 0 1 2 0 1 2 Magnetic Reid (T) Magnetic 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 temperatureTc (~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 ' m T= 1 5 0 K T= 2 3 0 K J j a - r '. ■ 3 • • • - • i - a • w ■ - .2 0 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 PZT Interface LSMO LSMO • Ti Zr 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 o o CM 20 30 40 50 Braggs 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 7 Sr0.33Mn0 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 ) /c A ~ r(+M)+n - m (61) 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 interface 10 10 0 P o 7.5 V -7.5 V MOKE signal (a.u) signal MOKE -10 V -10 V - 10-10 -1000 1000 -1000 1000 Magnetic 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 0.36 I (+M) 3.0.32 l(-M) 0.28 -1000 0 1000 H (Oe) Ug down 0.1 Ug up < 0.0 Ug (V) 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. 126 (a) Ug down Ug up -60 -15 0 15 Ug(V) (b) Ug down 0.1 Ug up 0.0 0 60-60 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. 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. 127 (a) 4 O O o CO -10 0 10 ug(V) (b) Ug down Ug up - + — -- 0.0 3 4 S (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 atz>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 - @ (z = 0) y LSMO / (? < 0) 1 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 Mx but also Mz 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. —- +10 V U ■ "1"1 * ■ — -10V 1 k t 136 ’ tz 0) ■M i 1.32 c g> ' Figure 6.8 P in S out MSHG 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. 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 d3z 2 _ r 2 orbital has lower energy than d x 2 _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 x 2 _ 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 t 2 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 x 2 _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 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 Lao 6 7 Cao3 3 Mn0 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 ±180 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 BIBLIOGRAPHY [1] M. Hosomi et al., A novel nonvolatile memory with spin torque transfer magnetization switching: spin-ram, IEDM tech. Dig. 459-462 (2005), DOI: 10.1109/IEDM.2005.1609379. [2] T. Kawahara et al., 2Mb spin-transfer torque RAM (SPRAM) with bit-by-bit bidirectional current write and parallelizing-direction current read, ISSCC Dig. Tech. Papers, 480-617 (2007), DOI: 10.1109/ISSCC.2007.373503. [3] C. Chappert, A. Fert and F. Nguyen Van Dau, The emergence of spin electronics in data storage, Nat. Mater. 6, 813-823 (2007), DOI: 10.1038/nmat2024. [4] D. Weller, A. Moser, L. Folks, M. E. Best, W. Lee, M. Toney, M. Schwickert, J.-U. Thiele, and M. Doerner, High Ku materials approach to 100 Gbits/in2, IEEE Trans. Magn. 36, 10-15 (2000), DOI: 10.1109/20.824418. [5] D. Alloyeau, C. Ricolleau, C. Mottet, T. Oikawa, C. Langlois, Y. Le Bouar, N. Braidy, and A. Loiseau, Size and shape effects on the order-disorder phase transition in CoPt nanoparticles, Nat. Mater. 8, 940-946 (2009), DOI:10.1038/nmat2574. [6] S. Sun, C. B. Murray, D. Weller, L. Folks, and A. Moser, Monodisperse FePt nanoparticles and ferromagnetic FePt nanocrystal superlattices, Science 287, 5460, 1989 (2000), DOI: 10.1126/science.287.5460.1989. [7] R. C. O’Handley, Modern Magnetic Materials (Wiley, New York, 2000), ISBN-13: 978-0471155669. [8] L. Neel, L'approche a la saturation de la magnetostriction, J. Phys. Radium 15, 376-378 (1954), DOI: 10.1051/jphysrad:01954001505037601. [9] A. D. Kent, Spintronics: Perpendicular all the way, Nat. Mater. 9, 699-700 (2010), DOI: 10.1038/nmat2844. [10] S. Ikeda, K. Miura, H. Yamamoto, K. Mizunuma, H. D. Gan, M. Endo, S. Kanai, J. Hayakawa, F. Matsukura, and H. Ohno, A perpendicular-anisotropy CoFeB-MgO magnetic tunnel junction, Nat. Mater. 9, 721-724 (2010), DOI: 10.1038/nmat2804. [11] W.-G. Wang, M. Li, S. Hageman, and C. L. Cbien, Electric-field-assisted switching in magnetic tunnel junctions, Nat. Mater. 11, 64-68 (2012), doi: 10.1038/nmat3171. [12] S. Mangin, D. Ravelosona, J. A. Katine, M. J. Carey, B. D. Terris, and E. E. Fullerton, Current-induced magnetization reversal in nanopillars with 139 perpendicular anisotropy, Nat. Mater. 5, 210-215 (2006), doi:10.1038/ nmat1595. [13] A. Brataas, A. D. Kent and H. Ohno, Current-induced torques in magnetic materials, Nat. Mater. 11, 372-381 (2012), doi:10.1038/nmat3311. [14] M. Bibes, A. Barthelemy, Multiferroics: Towards a magnetoelectric memory, Nature Mater. 7, 425-426 (2008), doi:10.1038/nmat2189. [15] Y.W. Yin, J. D. Burton, Y-M. Kim, A. Y. Borisevich, S. J. Pennycook, S. M. Yang, T.W. Noh, A. Gruverman, X. G. Li, E. Y. Tsymbal and Qi Li. Nat. Mater. 12, 397-402 (2013), doi:10.1038/nmat3564. [16] 14:25, 21 June 2005. Cassioli (26781 bytes) (Scheme of perpendicular recording technology. (Public Domain image by Luca Cassioli 2005)) [17] L. Landau and E. Lifshitz, On the Theory of the Dispersion of Magnetic Permeability in Ferromagnetic Bodies, Phys. Z. Sowjetunion 8, 153 (1935); T. L. Gilbert, A Lagrangian formulation of the gyromagnetic equation of the magnetic field, Phys. Rev. 100, 1243 (1955). [18] M. Gajek, M. Bibes, S. Fusil, K. Bouzehouane, J. Fontcuberta, A. Barthelemy & A. Fert, Tunnel junctions with multiferroic barriers, Nature Materials 6, 296 - 302 (2007), doi:10.1038/nmat1860. [19] John C. Mallinson, Magneto-Resistive and Spin Valve Heads: Fundamentals and Applications, ISBN 0-12466627-2. [20] Weinberger, P. (2008). "John Kerr and his Effects Found in 1877 and 1878". Philosophical Magazine Letters 88 (12): 897-907. Bibcode:2008PMagL..88..897W. doi: 10.1080/09500830802526604 [21] Kerr, John (1877). "On Rotation of the Plane of the Polarization by Reflection from the Pole of a Magnet". Philosophical Magazine 3: 321. doi: 10.1080/14786447708639245. [22] L. D. Landau and E. M. Lifshtz, Electrodynamics of Continuous Media (Pergamon, London, 1960). [23] Z. Q. Qiu, and S. D. Bader, Surface magneto-optic Kerr effect, Rev. Sci. Instrum. 71, 1243-1255 (2000), DOI: http://dx.doi.Org/10.1063/1.1150496. [24] S. Chikazumi, Physics of Ferromagnetism, vol. 94 of International Series of Monographs on Physics. Oxford, New York: Oxford Science Publications, 2nd ed„ 1997, ISBN-13: 978-0199564811. [25] M. van Kampen, C. Jozsa, J. T. Kohlhepp, P. LeClair, L. Lagae, W. J. M. de Jonge, and B. Koopmans, All-Optical Probe of Coherent Spin Waves, Phys. Rev. Lett. 88, 227201, (2002), DOI: http://dx.doi.org/10.1103/ PhysRevLett.88.227201. 140 [26] Th. Rasing, Nonlinear magneto-optics, J. Magn. Magn. Mater. 175, 35-50 (1997). [27] G. Luepke, Characterization of Semiconductor Interfaces by Second- Harmonic Generation, Surf. Sci. Reports 35, Nos. 3-4, 1999, pp. 75-161, DOI: 10.1016/S0167-5729(99)00007-2. [28] V. K. Valev, M. Gruyters, A. Kirilyuk, and Th. Rasing. Direct Observation of Exchange Bias Related Uncompensated Spins at the CoO/Cu Interface Phys. Rev. Lett. 96, 067206 (2006), DOI: http://dx.doi.org/10.1103/ PhysRevLett.96.067206. [29] Th. Gerrits, H. A. M. van den Berg, J. Hohlfeld, L. Bar, and Th. Rasing, Ultrafast precessional magnetization reversal by picosecond magnetic field pulse shaping. Nature (London) 418, 509 (2002), doi:10.1038/nature00905. [30] H. W. Schumacher, C. Chappert, R. C. Sousa, P. P. Freitas, and J. Miltat, Quasiballistic Magnetization Reversal, Phys. Rev. Lett. 90, 017204 (2003). DOI: http://dx.doi.Org/10.1103/PhysRevLett.90.017204. [31] S. I. Kiselev, J. C. Sankey, I.N. Krivorotov, N. C. Emley, R. J. Schoelkopf, R. A. Buhrman, and D. C. Ralph, Microwave oscillations of a nanomagnet driven by a spin-polarized current, Nature (London) 425, 380 (2003), doi: 10.1038/nature01967. [32] S. Kaka, M. R. Pufall, W. H. Rippard, T. J. Silva, S. E. Russek, and J. A. Katine, Mutual phase-locking of microwave spin torque nano-oscillators, Nature (London) 437, 389 (2005), doi:10.1038/nature04035. [33] P. Nemec, E. Rozkotova, N. Tesarova, F. Troj£nek, E. De Ranieri, K. Olejnlk, J. Zemen, V. Novak, M. Cukr, P. Maly, and T. Jungwirth, Experimental observation of the optical spin transfer torque, Nat. Phys. 8.411 (2012), doi: 10.1038/nphys2279. [34] B. Koopmans, J. J. M. Ruigrok, F. Dalla Longa, and W. J. M. de Jonge, Unifying Ultrafast Magnetization Dynamics, Phys. Rev. Lett. 95, 267207 (2005). DOI: http://dx.doi.org/10.1103/PhysRevLett.95.267207. [35] R. Urban, G. Woltersdorf, and B. Heinrich, Gilbert Damping in Single and Multilayer Ultrathin Films: Role of Interfaces in Nonlocal Spin Dynamics, Phys. Rev. Lett. 87, 217204 (2001), DOI: http://dx.doi.org/10.1103/ PhysRevLett.87.217204. [36] D. J. Twisselmann and R. D. McMichael, Intrinsic damping and intentional ferromagnetic resonance broadening in thin Permalloy films J. Appl. Phys. 93, 6903 (2003), DOI: http://dx.doi.Org/10.1063/1.1543884. [37] G. Woltersdorf, M. Buess, B. Heinrich, and C. H. Back, Time Resolved Magnetization Dynamics of Ultrathin Fe(001) Films: Spin-Pumping and Two- 141 Magnon Scattering, Phys. Rev. Lett. 95, 037401 (2005), DOI: http://dx.doi.org/10.1103/PhysRevLett.95.037401. [38] A. Barman, S. Wang, J. Maas, A. R. Hawkins, S. Kwon, J. Bokor, A. Liddle, and H. Schmidt, Size dependent damping in picosecond dynamics of single nanomagnets, Appl. Phys. Lett. 90, 202504 (2007), DOI: http://dx.doi.Org/10.1063/1.2740588. [39] A. Laraoui, J. Venuat, V. Halte, M. Albrecht, E. Beaurepaire, and J.-Y. Bigot, Study of individual ferromagnetic disks with femtosecond optical pulses, J. Appl. Phys. 101, 09C105 (2007), DOI: http://dx.doi.Org/10.1063/1.2711609. [40] S. Mizukami, Y. Ando, and T. Miyazaki, Effect of spin diffusion on Gilbert damping for a very thin permalloy layer in Cu/permalloy/Cu/Pt films, Phys. Rev. B 66, 104413 (2002), DOI: http://dx.doi.org/10.1103/ PhysRevB.66.104413. [41] Y. Fan, X. Ma, F. Fang, J. Zhu, Q. Li, T. P. Ma, Y. Z. Wu, Z. H. Chen, H. B. Zhao, and G. Luepke, Photoinduced spin angular momentum transfer into an antiferromagnetic insulator, Phys. Rev. B 89, 094428 (2014), DOI: http://dx.doi .org/10.1103/PhysRevB .89.094428. [42] V. Kambersky, Can. J. Phys. 48, 2906 (1970); J. Kunes and V. Kambersky, Phys. Rev. B 65, 212411 (2002). [43] B. Heinrich, D. Fraitova, and V. Kambersky, Phys. Status Solidi 23, 501 (1967); V. Kambersky, Czech. J. Phys. 26, 1366 (1976); V. Kambersky, First- principles investigation of the damping of fast magnetization precession in ferromagnetic 3d metals, Phys. Rev. B 76, 134416 (2007), DOI: http://dx.doi.Org/10.1103/PhysRevB.65.212411. [44] D. Steiauf and M. Fahnle, Damping of spin dynamics in nanostructures: An ab initio study, Phys. Rev. B 72, 064450 (2005), DOI: http://dx.doi.org/ 10.1103/PhysRevB.72.064450. [45] M. C. Hickey and J. S. Moodera, Origin of Intrinsic Gilbert Damping, Phys. Rev. Lett. 102, 137601 (2009), DOI: http://dx.doi.org/10.1103/PhysRevLett. 102.137601. [46] K. Gilmore, Y. U. Idzerda, and M. D. Stiles, dentification of the Dominant Precession-Damping Mechanism in Fe, Co, and Ni by First-Principles Calculations, Phys. Rev. Lett. 99, 027204 (2007), DOI: http://dx.doi.Org/10.1103/PhysRevLett.99.027204. [47] K. Gilmore, Y. U. Idzerda, and M. D. Stiles, Spin-orbit precession damping in transition metal ferromagnets, J. Appl. Phys. 103, 07D303 (2008). doi: 10.1063/1.2832348. 142 [48] A. Brataas, Y. Tserkovnyak, and G. E.W. Bauer, Scattering Theory of Gilbert Damping, Phys. Rev. Lett. 101, 037207 (2008). DOI: http://dx.doi.Org/10.1103/PhysRevLett. 101.037207. [49] H. Ebert, S. Mankovsky, D. Kodderitzsch, and P. J. Kelly, Ab Initio Calculation of the Gilbert Damping Parameter via the Linear Response Formalism, Phys. Rev. Lett. 107, 066603 (2011), DOI: http://dx.doi.Org/10.1103/PhysRevLett. 107.066603. [50]Y. Liu, A. A. Starikov, Z. Yuan, and P. J. Kelly, First-principles calculations of magnetization relaxation in pure Fe, Co, and Ni with frozen thermal lattice disorder, Phys. Rev. B 84, 014412 (2011), DOI: http://dx.doi.org/10.1103/ PhysRevB.84.014412. [51] J. Pelzl, R. Meckenstock, D. Spoddig, F. Schreiber, J. Pflaum and Z Frait, Spin-orbit-coupling effects on g-value and damping factor of the ferromagnetic resonance in Co and Fe films, J. Phys. Condens. Matter 15, S451 (2003), doi: 10.1088/0953-8984/15/5/302. [52] S. Ingvarsson, Gang Xiao, S. Parkin and R. Koch, Tunable magnetization damping in transition metal ternary alloys, Appl. Phys. Lett. 85, 4995 (2004), http://dx.doi.Org/10.1063/1.1828232. [53] J. O. Rantschler, R. D. McMichael.a), A. Castillo, A. J. Shapiro, W. F. Egelhoff Jr., B. B. Maranville, D. Pulugurtha, A. P. Chen and L. M. Connors, Effect of 3d, 4d, and 5d transition metal doping on damping in permalloy thin films J. Appl. Phys. 101, 033911 (2007), DOI: http://dx.doi.org/10.1063/ 1.2436471. [54] S. Mizukami, D. Watanabe, M. Oogane, Y. Ando, Y. Miura, M. Shirai and T. Miyazaki, Low damping constant for Co2FeAI Heusler alloy films and its correlation with density of states, J. Appl. Phys. 105, 07D306 (2009), DOI: http://dx.doi.Org/10.1063/1.3067607. [55] M. Oogane, T. Kubota, Y. Kota, S. Mizukami, H. Naganuma, A. Sakuma and Y. Ando, Gilbert magnetic damping constant of epitaxially grown Co based Heusler alloy thin films, Appl. Phys. Lett. 96, 252501 (2010), DOI: http://dx.doi.Org/10.1063/1.3456378. [56] H. Lee, Y.-H. A. Wang, C. K. A. Mewes, W. H. Butler, T. Mewes, S. Maat, B. York, M. J. Carey and J. R. Childress, Magnetization relaxation and structure of CoFeGe alloys, Appl. Phys. Lett. 95, 082502 (2009), DOI: http://dx.doi.Org/10.1063/1.3207749. [57] P. Bruno, Tight-binding approach to the orbital magnetic moment and magnetocrystalline anisotropy of transition-metal monolayers, Phys. Rev. B 39, 865 (1989), DOI: http://dx.doi.org/10.1103/PhysRevB.39.865. 143 [58] G. H. O. Daalderop, P. J. Kelly, and M. F. H. Schuurmans, Magnetic anisotropy of a free-standing Co monolayer and of multilayers which contain Co monolayers, Phys. Rev. B 50, 9989 (1994), DOI: http://dx.doi.org/10.1103/ PhysRevB.50.9989. [59] A. Matos-Abiague and J. Fabian, Anisotropic tunneling magneto- resistance and tunneling anisotropic magnetoresistance: Spin-orbit coupling in magnetic tunnel junctions, Phys. Rev. B 79, 155303 (2009), DOI: http://dx.doi.Org/10.1103/PhysRevB.79.155303. [60] C. Antoniak, J. Lindner, M. Spasova, D. Sudfeld, M. Acet, M. Farle, K. Fauth, U. Wiedwald, H.-G. Boyen, P. Ziemann, F. Wilhelm, A. Rogalev, and S. Sun, Enhanced Orbital Magnetism in Fe50Pt50 Nanoparticles, Phys. Rev. Lett. 97, 117201 (2006), DOI: http://dx.doi.org/10.1103/PhysRevLett.97. 117201. [61] Roman V. Chepulskii and W. H. Butler, Ab initio magnetocrystalline anisotropy at nanoscale: The case of FePt, Appl. Phys. Lett. 100, 142405 (2012), DOI: http://dx.doi.Org/10.1063/1.3700746. [62] H. Reichert, A. Schops, I. B. Ramsteiner, V. N. Bugaev, O. Shchyglo, A. Udyansky, H. Dosch, M. Asta, R. Drautz, and V. Honkimaki, Competition between Order and Phase Separation in Au-Ni, Phys. Rev. Lett. 95, 235703 (2005), DOI: http://dx.doi.org/10.1103/PhysRevLett.95.235703. [63] A. Cebollada, D. Weller, J. Sticht, G. R. Harp, R. F. C. Farrow, R. F. Marks, R. Savoy, and J. C. Scott, Enhanced magneto-optical Kerr effect in spontaneously ordered FePt alloys: Quantitative agreement between theory and experiment, Phys. Rev. B 50, 3419 (1994), DOI: http://dx.doi.org/ 10.1103/P hys RevB.50.3419. [64] C. Verdes, R. W. Chantrell, A. Satoh, J. W. Harrell, and D. Nikles, Self organisation, orientation and magnetic properties of FePt nanoparticle arrays, J. Magn. Magn. Mater. 304, 27 (2006), doi:10.1016/j.jmmm.2006.01.123. [65] P. Ravindran, A. Kjekshus, H. Fjellvag, P. James, L. Nordstrom, B. Johansson, and O. Eriksson, Large magnetocrystalline anisotropy in bilayer transition metal phases from first-principles full-potential calculations, Phys. Rev. B, 63, 144409 (2001), DOI: http://dx.doi.org/10.1103/ PhysRevB.63.144409. [66] I. V. Solovyev, P. H. Dederichs, and I. Mertig, Origin of orbital magnetization and magnetocrystalline anisotropy in TX ordered alloys (where T=Fe,Co and X=Pd,Pt), Phys. Rev. B, 52, 13419 (1995), DOI: http://dx.d 0 i.0 rg/l 0.1103/PhysRevB.52.13419. [67] Z. Lu, Roman V. Chepulskii, and W. H. Butler, First-principles study of magnetic properties of L10-ordered MnPt and FePt alloys, Phys. Rev. B, 81, 094437 (2010), DOI: http://dx.doi.org/10.1103/PhysRevB.81.094437. 144 [68] Gereon Meyer, and Jan-Ulrich Thiele, Effective electron-density dependence of the magnetocrystalline anisotropy in highly chemically ordered pseudobinary (Fe1-xMnx)50Pt50 L10 alloys, Phys. Rev. B 73, 214438 (2006), DOI: http://dx.doi.org/10-1103/PhysRevB.73.214438. [69] Till Burkert, Olle Eriksson, Sergei I. Simak, Andrei V. Ruban, Biplab Sanyal, Lars Nordstrom, and John M. Wills, Magnetic anisotropy of L10 FePt and Fe1-xMnxPt, Phys. Rev. B 71, 134411 (2005), DOI: http://dx.doi.org/ 10.1103/PhysRevB.71.134411. [70] S. Okamoto, N. Kikuchi, O. Kitakami, T. Miyazaki, Y. Shimada, and K. Fukamichi, Chemical-order-dependent magnetic anisotropy and exchange stiffness constant of FePt (001) epitaxial films, Phys. Rev. B 66, 024413 (2002), DOI: http://dx.doi.org/10.1103/PhysRevB.66.024413. [71] O. A. Ivanov, L.V. Solina, V. A. Demshina, and L. M. Magat, Fiz. Met. Metalloved. 35, 92 (1973). [72] G. Malinowski, K. C. Kuiper, R. Lavrijsen, H. J. M. Swagten, and B. Koopmans, Magnetization dynamics and Gilbert damping in ultrathin Co48Fe32B20 films with out-of-plane anisotropy, Appl. Phys. Lett. 94, 102501 (2009), DOI: http://dx.doi.org/10.1063/1-3093816. [73] P. He, X. Ma, J.W. Zhang, H. B. Zhao, G. Lupke, Z. Shi, and S. M. Zhou, Quadratic Scaling of Intrinsic Gilbert Damping with Spin-Orbital Coupling in L10 FePdPt Films: Experiments and Ab Initio Calculations, Phys. Rev. Lett. 110, 077203 (2013), DOI: http://dx.doi.org/10.1103/PhysRevLett.110.077203. [74] X. Ma, P. He, L. Ma, G. Y. Guo, H. B. Zhao, S. M. Zhou, and G. Liipke, Spin-orbit interaction tuning of perpendicular magnetic anisotropy in L10 FePdPt films, Appl. Phys. Lett. 104, 192402 (2014), DOI: http://dx.doi.org/ 10.1063/1.4876128. [75] J.-P. Kuang, M. Kontani, M. Matsui, and K. Adachi, Physica B 149, 209 (1988). [76] A. Alam, B. Kraczek, and D. D. Johnson, Structural, magnetic, and defect properties of Co-Pt-type magnetic-storage alloys: Density-functional theory study of thermal processing effects, Phys. Rev. B 82, 024435 (2010), DOI: http://dx.doi.org/10.1103/PhysRevB.82.024435. [77] C. J. Aas, L. Szunyogh, J. S. Chen, and R. W. Chantrell, Magnetic anisotropy of FePt: Effect of lattice distortion and chemical disorder, App. Phys. Lett. 99, 132501 (2011), DOI: http://dx.doi.Org/10.1063/1.3644478. [78] P. Kamp, A. Marty, B. Gilles, R. Hoffmann, S. Marchesini, M. Belakhovsky, C. Boeglin, H. Durr, S. Dhesi, and G. van der Laan and A. Rogalev, Correlation of spin and orbital anisotropies with chemical order in 145 Fe0.5Pd0.5 alloy films using magnetic circular x-ray dichroism, Phys. Rev. B 59, 1105 (1999), DOI: http://dx.doi.org/10.1103/PhysRevB.59.1105. [79] S. Dhesi, G. van der Laan, H.A. Durr, M. Belakhovsky, S. Marchesini, P. Kamp, A. Marty, B. Gilles, A. Rogalev, Magnetocrystalline anisotropy in FePd alloys studied using transverse X-ray magnetic circular dichroism, J. Magn. Magn. Mat. 226-230 (2001) 1580-1582, doi:10.1016/S0304-8853(00)00904-5. [80] J. P. Nibarger, R. Lopusnik, Z. Celinski, and T. J. Silva, Variation of magnetization and the Lande g factor with thickness in Ni-Fe films, App. Phys. Lett. 83, 93 (2003), DOI: http://dx.doi.Org/10.1063/1.1588734. [81] J. Lyubina, I. Opahle, M. Richter, O. Gutfleisch, K. Muller, and L. Schultz, Influence of composition and order on the magnetism of Fe-Pt alloys: Neutron powder diffraction and theory , App. Phys. Lett. 89, 032506 (2006), DOI: http://dx.doi.Org/10.1063/1.2222244. [82] M. Chen, Z. Shi, W. J. Xu, X. X. Zhang, J. Du, and S. M. Zhou, Tuning anomalous Hall conductivity in L10 FePt films by long range chemical ordering, App. Phys. Lett. 98, 082503 (2011), DOI: http://dx.doi.org/ 10.1063/1.3556616. [83] C. Stamm, T. Kachel, N. Pontius, R. Mitzner, T. Quast, K. Holldack, S. Khan, C. Lupulescu, E. F. Aziz, M. Wietstruk, H. A. Durr and W. Eberhardt, Femtosecond modification of electron localization and transfer of angular momentum in nickel, Nature Mater. 6, 740 - 743 (2007), doi: 10.1038/ nmat1985. [84] B. Koopmans, G. Malinowski, F. Dalla Longa, D. Steiauf, M. Fahnle, T. Roth, M. Cinchetti and M. Aeschlimann, Explaining the paradoxical diversity of ultrafast laser-induced demagnetization, Nature Mater. 9, 259-265 (2010), doi: 10.1038/nmat2593. [85] Dennis Rudolf, Chan La-O-Vorakiat, Marco Battiato, Roman Adam, Justin M. Shaw, Emrah Turgut, Pablo Maldonado, Stefan Mathias, Patrik Grychtol, Hans T. Nembach, Thomas J. Silva, Martin Aeschlimann, Henry C. Kapteyn, Margaret M. Murnane, Claus M. Schneider and Peter M. Oppeneer, Ultrafast magnetization enhancement in metallic multilayers driven by superdiffusive spin current, Nature Commun. 3, 1037 (2012), doi:10.1038/ ncomms2029. [86] Gyung-Min Choi, Byoung-Chul Min, Kyung-Jin Lee and David G. Cahill, Spin current generated by thermally driven ultrafast demagnetization, Nature Commun. 5, 4334 (2014), doi:10.1038/ncomms5334. [87] P. Nemec, E. Rozkotova, N. Tesarova, F. Trojanek, E. De Ranieri, K. Olejnik, J. Zemen, V. Novak, M. Cukr, P. Maly and T. Jungwirth, Experimental 146 observation of the optical spin transfer torque, Nature Phys. 8, 411-415 (2012), doi: 10.1038/nphys2279. [88] N. Tesarova, P. Nemec, E. Rozkotova, J. Zemen, T. Janda, D. Butkovicova, F. Troja'nek, K. Olejnik, V. Novak, P. Maly and T. Jungwirth, Experimental observation of the optical spin-orbit torque, Nature Photo. 7, 492 (2013), doi: 10.1038/nphoton.2013.76. [89] Natsuki Kanda, Takuya Higuchi, Hirokatsu Shimizu, Kuniaki Konishi, Kosuke Yoshioka and Makoto Kuwata-Gonokami, The vectorial control of magnetization by light, Nature Commun. 2, 362 (2011), doi:10.1038/ ncomms1366. [90] A. R. Mellnik, J. S. Lee, A. Richardella, J. L. Grab, P. J. Mintun, M. H. Fischer, A. Vaezi, A. Manchon, E.-A. Kim, N. Samarth and D. C. Ralph, Spin- transfer torque generated by a topological insulator, Nature(London) 511, 449 (2014), doi:10.1038/nature13534. [91] I. N. Krivorotov, N. C. Emley, J. C. Sankey, S. I. Kiselev, D. C. Ralph, R. A. Buhrman, Time-Domain Measurements of Nanomagnet Dynamics Driven by Spin-Transfer Torques, Science 307, 228 (2005), DOI:10.1126/ science. 1105722. [92] L. Liu, C. F. Pai, D. C. Ralph, and R. A. Buhrman, Magnetic Oscillations Driven by the Spin Hall Effect in 3-Terminal Magnetic Tunnel Junction Devices, Phys. Rev. Lett. 109, 186602 (2012), DOI: http://dx.doi.org/ 10.1103/PhysRevLett. 109.186602. [93] Q. Zhang, A. V. Nurmikko, A. Anguelouch, G. Xiao, and A. Gupta, Coherent Magnetization Rotation and Phase Control by Ultrashort Optical Pulses in Cr02 Thin Films, Phys. Rev. Lett. 89, 177402 (2002), DOI: http://dx.doi.Org/10.1103/PhysRevLett.89.177402. [94] S. Tomimoto, M. Matsubara, T. Ogasawara, H. Okamoto, T. Kimura, and Y. Tokura, Optical Control of the Magnetic Anisotropy of Ferromagnetic Bilayered Manganites, Phys. Rev. Lett. 98, 017402 (2007), DOI: http://dx.doi.Org/10.1103/PhysRevLett.98.017402. [95] A. V. Scherbakov, A. S. Salasyuk, A. V. Akimov, X. Liu, M. Bombeck, C. Bruggemann, D. R. Yakovlev, V. F. Sapega, J. K. Furdyna, and M. Bayer, Coherent Magnetization Precession in Ferromagnetic (Ga,Mn)As Induced by Picosecond Acoustic Pulses, Phys. Rev. Lett. 105, 117204 (2010), DOI: http://dx.doi.Org/10.1103/PhysRevLett. 105.117204. [96] G. Ju, A. V. Nurmikko, R. F. C. Farrow, R. F. Marks, M. J. Carey, and B. A. Gurney, Ultrafast Time Resolved Photoinduced Magnetization Rotation in a Ferromagnetic/Antiferromagnetic Exchange Coupled System, Phys. Rev. Lett. 82, 3705 (1999), DOI: http://dx.doi.org/10.1103/PhysRevLett.82.3705. 147 [97] H. B. Zhao, D. Talbayev, X. Ma, Y. H. Ren, A. Venimadhav, Q. Li, and G. Lupke, Coherent Spin Precession via Photoinduced Antiferromagnetic Interactions in La0.67Ca0.33MnO3, Phys. Rev. Lett. 107, 207205 (2011), DOI: http://dx.doi .org/10.1103/PhysRevLett. 107.207205. [98] Y. Hashimoto, S. Kobayashi, and H. Munekata, Photoinduced Precession of Magnetization in Ferromagnetic (Ga,Mn)As, Phys. Rev. Lett. 100, 067202 (2008), DOI: http://dx.doi.org/10.1103/PhysRevLett.100.067202. [99] J. Qi, Y. Xu, A. Steigerwald, X. Liu, J. K. Furdyna, I. E. Perakis, and N. H. Tolk, Ultrafast laser-induced coherent spin dynamics in ferromagnetic Ga1-xMnxAs/GaAs structures, Phys. Rev. B 79, 085304 (2009), DOI: http://dx.doi.org/10.1103/PhysRevB.79.085304. [100] D. M. Wang, Y. H. Ren, X. Liu, J. K. Furdyna, M. Grimsditch, and R. Merlin, Light-induced magnetic precession in (Ga.Mn)As slabs: Hybrid standing-wave Damon-Eshbach modes, Phys. Rev. B 75, 233308 (2007), DOI: http://dx.doi.Org/10.1103/PhysRevB.75.233308. [101] H. Yuan, H. Gao, Y. Gong, J. Lu, X. Zhang, J. Zhao, Y. H. Ren, H. B. Zhao, and L. Chen, Photoinduced Spin Precession in Fe/GaAs(001) Heterostructure with Low Power Excitation, Appl. Phys. Expr. 6, 073008 (2013), doi: 10.7567/APEX.6.073008. [102] W. N. Cao, J. Li, G. Chen, J. Zhu, C. R. Hu, and Y. Z. Wu, Temperature-dependent magnetic anisotropies in epitaxial Fe/CoO/MgO(001) system studied by the planar Hall effect, Appl. Phys. Lett. 98, 262506 (2011), DOI: http://dx.doi.Org/10.1063/1.3606531. [103] Y. Fan, X. Ma, F. Fang, J. Zhu, Q. Li, T. P. Ma, Y. Z. Wu, Z. H. Chen, H. B. Zhao, and G. Luepke, Photoinduced spin angular momentum transfer into an antiferromagnetic insulator, Phys. Rev. B 89, 094428 (2014), DOI: http://dx.doi.org/10.1103/PhysRevB.89.094428. [104] A. V. Kimel, A. Kirilyuk, P. A. Usachev, R. V. Pisarev, A. M. Balbashov and Th. Rasing, Ultrafast non-thermal control of magnetization by instantaneous photomagnetic pulses, Nature (London) 435, 655 (2005), doi: 10.1038/nature03564. [105] A. V. Kimel, A. Kirilyuk, A. Tsvetkov, R. V. Pisarev and Th. Rasing, Laser-induced ultrafast spin reorientation in the antiferromagnet Tm Fe03, Nature (London) 429, 850 (2004), doi:10.1038/nature02659. [106] Ch. Kant, T. Rudolf, F. Schrettle, F. Mayr, J. Deisenhofer, P. Lunkenheimer, M. V. Eremin, and A. Loidl, Optical spectroscopy in CoO: Phononic, electric, and magnetic excitation spectrum within the charge- transfer gap, Phys. Rev. B 78, 245103 (2008), DOI: http://dx.doi.org/ 10.1103/PhysRevB.78.245103. 148 [107] G. M. Muller, J. Walowski, M. Djordjevic, G.-X. Miao, A. Gupta, A. V. Ramos, K. Gehrke, V. Moshnyaga, K. Samwer, J. Schmalhorst, A. Thomas, A. Hutten, G. Reiss, J. S. Moodera and M. Munzenberg, Spin polarization in half-metals probed by femtosecond spin excitation, Nature Mater. 8, 56 (2009), doi:10.1038/nmat2341. [108] E. Beaurepaire, J.-C. Merle, A. Daunois, and J.-Y. Bigot, Ultrafast Spin Dynamics in Ferromagnetic Nickel, Phys. Rev. Lett. 76, 4250 (1996), DOI: http://dx.doi.org/10.1103/PhysRevLett.76.4250. [109] J. O. Rantschler, R. D. McMichael,a), A. Castillo, A. J. Shapiro, W. F. Egelhoff Jr., B. B. Maranville, D. Pulugurtha, A. P. Chen and L. M. Connors, Effect of 3d, 4d, and 5d transition metal doping on damping in permalloy thin films, J. Appl. Phys. 101, 033911 (2007), http://dx.doi.Org/10.1063/1.2436471. [110] C. M. Jiang, L. R. Baker, J. M. Lucas, J. V.-Weis, A. P. Alivisatos, and S. R. Leone, Characterization of Photo-Induced Charge Transfer and Hot Carrier Relaxation Pathways in Spinel Cobalt Oxide (Co304), J. Phys. Chem. C DOI: 10.1021/jp5071133. [111] H. X. Deng, J. Li, S. S. Li, J. B. Xia, A. Walsh, and S. H. Wei, Origin of antiferromagnetism in CoO: A density functional theory study, Appl. Phys. Lett 96, 162508 (2010), http://dx.doi.Org/10.1063/1.3402772. [112] V. K. Valev, M. Gruyters, A. Kirilyuk, and Th. Rasing, Direct Observation of Exchange Bias Related Uncompensated Spins at the CoO/Cu Interface, Phys. Rev. Lett. 96, 067206 (2006), DOI: http://dx.doi.org/10.1103/ PhysRevLett.96.067206. [113] Y. Liu, L. R. Shelford, V. V. Kruglyak, R. J. Hicken, Y. Sakuraba, M. Oogane, and Y. Ando, Optically induced magnetization dynamics and variation of damping parameter in epitaxial Co2MnSi Heusler alloy films, Phys. Rev. B 81, 094402 (2010), DOI: http://dx.doi.org/10.1103/ PhysRevB.81.094402. [114] Fredrik Hansteen, Alexey Kimel, Andrei Kirilyuk, and Theo Rasing, Femtosecond Photomagnetic Switching of Spins in Ferrimagnetic Garnet Films, Phys. Rev. Lett. 95, 047402 (2005), DOI: http://dx.doi.org/10.1103/ PhysRevLett.95.047402. [115] Andrei Kirilyuk, Alexey V. Kimel, and Theo Rasing, Ultrafast optical manipulation of magnetic order, Rev. Mod. Phys. 82, 2731 (2010), DOI: http://dx.doi.Org/10.1103/RevModPhys.82.2731. [116] V. V. Volkov, Z. L. Wang, B. S. Zou, Carrier recombination in clusters of NiO, Chem. Phys. Lett. 337, 117 (2001), doi: 10.1016/S0009-2614(01)00191- 9. 149 [117] H.W. Schumacher et al., Phase Coherent Precessional Magnetization Reversal in Microscopic Spin Valve Elements, Phys. Rev. Lett. 90, 017201 (2003), DOI: http://dx.doi.org/10.1103/PhysRevLett.90.017201. [118] M. van Kampen et al., All-Optical Probe of Coherent Spin Waves, Phys. Rev. Lett. 88, 227201 (2002), DOI: http://dx.doi.org/10.1103/ PhysRevLett.88.227201. [119] H. B. Zhao et al., Ultrafast magnetization dynamics of epitaxial Fe films on AIGaAs (001), Appl. Phys. Lett. 86, 152512 (2005), DOI: http://dx.doi.Org/10.1063/1.1900939. [120] M. Djordjevic and M. Munzenberg, Connecting the timescales in picosecond remagnetization experiments, Phys. Rev. B 75, 012404 (2007), DOI: http://dx.doi.Org/10.1103/PhysRevB.75.012404. [121] C. D. Stanciu et al., Ultrafast Interaction of the Angular Momentum of Photons with Spins in the Metallic Amorphous Alloy GdFeCo, Phys. Rev. Lett. 98,207401 (2007), DOI: http://dx.doi.org/10.1103/PhysRevLett.98.207401. [122] D. Talbayev et al., Photoinduced coherent magnetization precession in epitaxial La0.67Ca0.33MnO3 films, Phys. Rev. B 73, 014417 (2006), DOI: http://dx.doi.Org/10.1103/PhysRevB.73.014417. [123] T. Ogasawara et al., Photoinduced spin dynamics in La0.6Sr0.4MnO3 observed by time-resolved magneto-optical Kerr spectroscopy, Phys. Rev. B 68, 180407(R) (2003), DOI: http://dx.doi.org/10.1103/PhysRevB.68.180407. [124] R. D. Averitt, A. I. Lobad, C. Kwon, S. A. Trugman, V. K. Thorsmolle, and A. J. Taylor, Ultrafast Conductivity Dynamics in Colossal Magneto- resistance Manganites, Phys. Rev. Lett. 87, 017401 (2001), DOI: http://dx.doi.Org/10.1103/Phys RevLett.87.017401. [125] H. B. Zhao, D. Talbayev, Y. Fan, G. Lupke, A. T. Hanbicki, C. H. Li and B. T. Jonker, Photoinduced spin waves in Fe/AIGaAs (001) heterostructure, Phys. Status Solidi C 5, 2627 (2008), DOI: 10.1002/ pssc.200779107. [126] B. Lenk, H. Ulrichs, F. Garbs, M. Munzenberg, The building blocks of magnonics, Phys. Rep. 507, 107 (2011), doi: 10.1016/j.physrep.2011.06.003. [127] J Walowski, M Djordjevic Kaufmann, B Lenk, C Hamann, J McCord and M Munzenberg, Intrinsic and non-local Gilbert damping in polycrystalline nickel studied by Ti: sapphire laser fs spectroscopy, J. Phys. D 41, 164016 (2008), doi: 10.1088/0022-3727/41/16/164016. [128] G. M. Muller, M. Munzenberg, G.-X. Miao, and A. Gupta, Activation of additional energy dissipation processes in the magnetization dynamics of epitaxial chromium dioxide films, Phys. Rev. B 77, 020412(R) (2008), DOI: http://dx.doi.Org/10.1103/PhysRevB.77.020412. 150 [129] M. Uehara, S. Mori, C. H. Chen and S.-W. Cheong, Percolative phase separation underlies colossal magnetoresistance in mixed-valent manganites , Nature (London) 399, 560 (1999), doi:10.1038/21142. [130] M. Fath, S. Freisem, A. A. Menovsky, Y. Tomioka, J. Aarts, J. A. Mydosh, Spatially Inhomogeneous Metal-lnsulator Transition in Doped Manganites, Science 285, 1540 (1999), DOI: 10.1126/science.285.5433.1540. [131] Y. Tokura, Critical features of colossal magnetoresistive manganites, Rep. Prog. Phys. 69, 797 (2006), doi:10.1088/0034-4885/69/3/R06. [132] K. H. Ahn, T. Lookman and A. R. Bishop, Strain-induced metal-insulator phase coexistence in perovskite manganites, Nature (London) 428, 401 (2004), doi: 10.1038/nature02364. [133] Amlan Biswas, M. Rajeswari, R. C. Srivastava, T. Venkatesan, R. L. Greene, Q. Lu, A. L. de Lozanne, and A. J. Millis, Strain-driven charge- ordered state in La0.67Ca0.33MnO3, Phys. Rev. B 63, 184424 (2001), DOI: http://dx.doi.Org/10.1103/PhysRevB.63.184424. [134] R. Rauer, J. Backstrom, D. Budelmann, M. KurfiB, M. Schilling, M. Rubhausen, T. Walter, K. Dorr and S. L. Cooper, Thickness dependent phase separation in La0.7Ca0.3MnO3 films, Appl. Phys. Lett. 81, 3777 (2002), DOI: http://dx.doi.Org/10.1063/1.1520705. [135] V. Moshnyaga, A. Giske, K. Samwer, E. Mishina, T. Tamura, S. Nakabayashi, A. Belenchuk, O. Shapoval and L. Kulyuk, Giant negative photoconductivity in La0.7Ca0.3MnO3 thin films, J. Appl. Phys. 95, 7360 (2004), DOI: http://dx.doi.Org/10.1063/1.1687555. [136] T. Ogasawara, K. Ohgushi, Y. Tomioka, K. S. Takahashi, H. Okamoto, M. Kawasaki, and Y. Tokura, General Features of Photoinduced Spin Dynamics in Ferromagnetic and Ferrimagnetic Compounds, Phys. Rev. Lett. 94, 087202 (2005), DOI: http://dx.doi.org/10.1103/PhysRevLett.94.087202. [137] Colossal Magnetoresistive Oxides, edited by Y. Tokura (Gordon and Breach, London, 1999), ISBN-13: 978-9056992316. [138] M. Matsubara, Y. Okimoto, T. Ogasawara, Y. Tomioka, H. Okamoto, and Y. Tokura, Ultrafast Photoinduced Insulator-Ferromagnet Transition in the Perovskite Manganite Gd0.55Sr0.45MnO3, Phys. Rev. Lett. 99, 207401 (2007), DOI: http://dx.doi.org/10.1103/PhysRevLett.99.207401. [139] M. Quijada, J. Cerne, J. R. Simpson, H. D. Drew, K. H. Ahn, A. J. Millis, R. Shreekala, R. Ramesh, M. Rajeswari, and T. Venkatesan, Optical conductivity of manganites: Crossover from Jahn-Teller small polaron to coherent transport in the ferromagnetic state, Phys. Rev. B 58, 16093 (1998), DOI: http://dx.doi.org/10.1103/PhysRevB.58.16093. 151 [140] N. N. Kovaleva, A. V. Boris, C. Bernhard, A. Kulakov, A. Pimenov, A. M. Balbashov, G. Khaliullin, and B. Keimer, Spin-Controlled Mott-Hubbard Bands in LaMn03 Probed by Optical Ellipsometry, Phys. Rev. Lett. 93, 147204 (2004), DOI: http://dx.doi.org/10.1103/PhysRevLett.93.147204. [141] A. Ohtomo, and H. Y. Hwang. A high-mobility electron gas at the LaAI03/SrTi03 heterointerface, Nature 427, 423-426 (2004), doi:10.1038/ nature02308. [142] K. Ueda, H. Tabata, and T. Kawai, Ferromagnetism in LaFe03-LaCr03 Superlattices, Science 280, 1064-1066 (1998), DOL10.1126/ science.280.5366.1064. [153] J. Chakhalian, J. W. Freeland, H.-U. Habermeier, G. Cristiani, G. Khaliullin, M. van Veenendaal, and B. Keimer. Orbital Reconstruction and Covalent Bonding at an Oxide Interface, Science 318, 1114-1117 (2007), DOI: 10.1126/science. 1149338. [144] Y. Tokura and H. Y. Hwang. Condensed-matter physics: Complex oxides on fire, Nature Mater. 7, 694 (2008), doi:10.1038/nmat2264. [145] J. Mannhart and D.G. Schlom, Oxide Interfaces— An Opportunity for Electronics, Science 327,1607-1611 (2010), DOL10.1126/science.1181862. [146] N. Reyren, S. Thiel, A. D. Caviglia, L. Fitting Kourkoutis, G. Hammerl, C. Richter, C. W. Schneider, T. Kopp, A.-S. Ruetschi, D. Jaccard, M. Gabay, D. A. Muller, J.-M. Triscone, and J. Mannhart, Superconducting Interfaces Between Insulating Oxides, Science 317, 1196-1199 (2007), DOI: 10.1126/ science. 1146006. [147] V. Garcia, M. Bibes, L. Bocher, S. Valencia, F. Kronast, A. Crassous, X. Moya, S. Enouz-Vedrenne, A. Gloter, D. Imhoff, C. Deranlot, N. D. Mathur, S. Fusil, K. Bouzehouane, and A. Barthelemy, Ferroelectric Control of Spin Polarization, Science 327, 1106-1110 (2010), DOI: 10.1126/science. 1184028. [148] Z. Fang, I.V. Solovyev, and K. Terakura, Phase Diagram of Tetragonal Manganites, Phys. Rev. Lett. 84, 3169 (2000), DOI: http://dx.doi.org/10.1103/ PhysRevLett.84.3169. [149] C. A. F. Vaz, J. Hoffman, Y. Segal, J. W. Reiner, R. D. Grober, Z. Zhang, C. H. Ahn, and F. J. Walker, Origin of the Magnetoelectric Coupling Effect in Pb(Zr0.2Ti0.8)03/La0.8Sr0.2Mn03 Multiferroic Heterostructures, Phys. Rev. Lett 104, 127202 (2010), DOI: http://dx.doi.org/10.1103/PhysRevLett. 104.127202; Hajo J. A. Molegraaf, Jason Hoffman, Carlos A. F. Vaz, Stefano Gariglio, Dirk van der Marel, Charles H. Ahn, Jean-Marc Triscone, Magnetoelectric Effects in Complex Oxides with Competing Ground States, Advanced Materials 21, 3470 (2009), DOI: 10.1002/adma.200900278. 152 [150] Shuai Dong, Xiaotian Zhang, Rong Yu, J.-M. Liu, and Elbio Dagotto, Microscopic model for the ferroelectric field effect in oxide heterostructures, Phys. Rev. B 84, 155117 (2011), DOI: http://dx.doi.org/10.1103/PhysRevB. 84.155117. [151] J. D. Burton and E.Y. Tsymbal, Giant Tunneling Electroresistance Effect Driven by an Electrically Controlled Spin Valve at a Complex Oxide Interface, Phys. Rev. Lett. 106, 157203 (2011), DOI: http://dx.doi.org/10.1103/ PhysRevLett. 106.157203. [152] J. Garcia-Barriocanal, J. C. Cezar, F. Y. Bruno, P. Thakur, N. B. Brookes, C. Utfeld, A. Rivera-Calzada, S. R. Giblin, J. W. Taylor, J. A. Duffy,S .B. Dugdale, T. Nakamura, K. Kodama, C. Leon, S. Okamoto, and J. Santamaria, Spin and orbital Ti magnetism at LaMn03/SrTi03 interfaces, Nature Commun. 1, 82 (2010), doi:10.1038/ncomms1080. [153] Y. Tokura, and N. Nagaosa, Orbital Physics in Transition-Metal Oxides, Science 288, 462-468 (2000), DOL10.1126/science.288.5465.462. [154] I. V. Solovyev and K. Terakura, Zone Boundary Softening of the Spin- Wave Dispersion in Doped Ferromagnetic Manganites, Phys. Rev. Lett. 82, 2959-2962 (1999), DOI: http://dx.doi.org/10.1103/PhysRevLett.82.2959. [155] P. -G. de Gennes, Effects of Double Exchange in Magnetic Crystals, Phys. Rev. 118, 141-154 (1960), DOI: http://dx.doi.org/10.1103/ PhysRev.118.141. [156] M. Matsuda, R. S. Fishman, T. Hong, C. H. Lee, T. Ushiyama, Y. Yanagisawa, Y. Tomioka, and T. Ito, Magnetic Dispersion and Anisotropy in Multiferroic BiFe03, Phys. Rev. Lett. 109, 067205 (2012), DOI: http://dx.doi.Org/10.1103/PhysRevLett. 109.067205. [157] Jaehong Jeong, E. A. Goremychkin, T. Guidi, K. Nakajima, Gun Sang Jeon, Shin-Ae Kim, S. Furukawa. Spin Wave Measurements over the Full Brillouin Zone of Multiferroic BiFe03, Phys. Rev. Lett. 108, 077202 (2012), DOI: http://dx.doi.Org/10.1103/PhysRevLett. 108.077202. [158] John B. Goodenough. Magnetism and the chemical bond (Interscience- Wiley, New York, 1963) [159] Isaac B. Bersuker, Pseudo Jahn-Teller Origin of Perovskite Multiferroics, Magnetic-Ferroelectric Crossover, and Magnetoelectric Effects: The d0-d10 Problem, Phys. Rev. Lett. 108, 137202 (2012), DOI: http://dx.doi.org/10.1103/ PhysRevLett. 108.137202. [160] H. B. Zhao, D. Talbayev, X. Ma, Y. H. Ren, A. Venimadhav, Qi Li, and G. Lupke. Coherent Spin Precession via Photoinduced Antiferromagnetic Interactions in La0.67Ca0.33MnO3, Phys. Rev. Lett. 107, 207205 (2011), DOI: http://dx.doi.Org/10.1103/PhysRevLett. 107.207205. 153 [161] J. F. Scott and C. A. Paz de Araujo, Ferroelectric Memories, Science 246, 1400 (1989), D0l:10.1126/science.246.4936.1400. [162] Sandra Dussan, Ashok Kumar, J. F. Scott, and Ram S. Katiyar, Magnetic effects on dielectric and polarization behavior of multiferroic heterostructures, Appl. Phys. Lett. 96, 072904 (2010), DOI: http://dx.doi.org/ 10.1063/1.3327889. [163] J. H. Park, E. Vescoso, H. J. Kim, C. Kwon, R. Ramesh, and T. Venkatesan, Direct evidence for a half-metallic ferromagnet, Nature (London) 392, 794-796 (1998), doi: 10.1038/33883. [164] M. Bowen, M. Bibes, A. Barthelemy, J.-P. Contour, A. Anane, Y. LemaTtre, and A. Fert. Nearly total spin polarization in La2/3Sr1/3Mn03 from tunneling experiments, Appl. Phys. Lett. 82, 233 (2003), http://dx.doi.org/ 10.1063/1.1534619. [165] Hiroyuki Yamada, Yoshihiro Ogawa, Yuji Ishii, Hiroshi Sato, Masashi Kawasaki, Hiroshi Akoh, Yoshinori Tokura, Engineered Interface of Magnetic Oxides, Science 305, 646 (2004), DOI: 10.1126/science. 1098867. [166] Y. Fan, K. J. Smith, G. Lupke, A. T. Hanbicki, R. Goswami, C. H. Li, H. B. Zhao and B. T. Jonker, Exchange bias of the interface spin system at the Fe/MgO interface, Nature Nanotechnology 8, 438-444 (2013), doi:10.1038/ nnano.2013.94. [167] H. B. Zhao, D. Talbayev, G. Liipke, A. T. Hanbicki, C. H. Li, O. M. J. vant Erve, G. Kioseoglou, B. T. Jonker, Interface Magnetization Reversal and Anisotropy in Fe/AIGaAs(001), Phys. Rev. Lett. 95, 2005, pp. 137202-5, DOI: http://dx.doi.Org/10.1103/PhysRevLett.95.137202 [168] J. Dho, N. H. Hur, I. S. Kim, and Y. K. Park, Oxygen pressure and thickness dependent lattice strain in La0.7Sr0.3MnO3 films, J. Appl. Phys. 94, 7670-4 (2003), DOI: http://dx.doi.Org/10.1063/1.1628831. [169] C. Aruta, G. Balestrino, A. Tebano, G. Ghiringhelli and N. B. Brookes, Cooperative enhancement of in-plane orbital ordering by oxygen deficiency and in-plane tensile strain in La0.7Sr0.3MnO3-6 thin films, Euro-Phys. Lett. 80 (2007) 37003, doi: 10.1209/0295-5075/80/37003. [170] A. Q. Jiang, Z. H. Chen, W. Y. Hui, D. Wu, and J. F. Scott, Subpicosecond Domain Switching in Discrete Regions of Pb(Zr0.35Ti0.65)03 Thick Films, Adv. Funct. Mater. 2012, 22, 2148-2153, DOI: 10.1002/adfm.201102829. [1 7 1 ] P. Gerber, U. Bottger, and R. Waser, Composition influences on the electrical and electromechanical properties of lead zirconate titanate thin films, J. Appl. Phys. 100, 124105 (2006), DOI: http://dx.doi.Org/10.1063/1.2401047. 154 [172] J. D. Burton and E. Y. Tsymbal, Evolution of the band alignment at polar oxide interfaces, Phys. Rev. B 82, 161407 (R) (2010), DOI: http://dx.doi.org/10.1103/PhysRevB.82.161407. [173] A. Tebano, C. Aruta, S. Sanna, P. G. Medaglia, and G. Balestrino, Evidence of Orbital Reconstruction at Interfaces in Ultrathin La0.67Sr0.33MnO3 Films, Phys. Rev. Lett. 100, 137401 (2008), DOI: http://dx.doi.Org/10.1103/PhysRevLett. 100.137401. [174] A. Tebano, A. Orsini, P. G. Medaglia, D. Di Castro, and G. Balestrino, B. Freelon, A. Bostwick, Young Jun Chang, G. Gaines, and E. Rotenberg, Preferential occupation of interface bands in La2/3Sr1/3Mn03 films as seen via angle-resolved photoemission, Phys. Rev. B 82, 214407 (2010), DOI: http://dx.doi.Org/10.1103/PhysRevB.82.214407. [175] C. Aruta, G. Ghiringhelli, V. Bisogni, L. Braicovich, N. B. Brookes, A. Tebano, and G. Balestrino, Orbital occupation, atomic moments, and magnetic ordering at interfaces of manganite thin films, Phys. Rev. B 80, 014431 (2009), DOI: http://dx.doi.org/10.1103/PhysRevB.80.014431. [176] D. Pesquera, G. Herranz, A. Barla, E. Pellegrin, F.Bondino, E. Magnano, F. Sanchez and J. Fontcuberta. Surface symmetry-breaking and strain effects on orbital occupancy in transition metal perovskite epitaxial films, Nature Commun. 3. 1189(2012), doi:10.1038/ncomms2189. 155 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. 156