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2015
Optical characterization of ferromagnetic and multiferroic thin-Optical characterization of ferromagnetic and multiferroic thin-
film heterostructures film heterostructures
Xin Ma College of William & Mary - Arts & Sciences
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Optical Characterization of Ferromagnetic and Multiferroic Thin-Film Heterostructures
Xin Ma
Lingbi, Suzhou, Anhui, P. R. China
Bachelor of Science, the University of Science and Technologyof China, 2009
A Dissertation presented to the Graduate Faculty of the College of William and Mary in Candidacy for the Degree of
Doctor of Philosophy
Department of Applied Science
The College of William and Mary January, 2015
APPROVAL PAGE
This Dissertation is submitted in partial fulfillment of the requirements for the degree of
Doctor of Philosophy
10 I ' I o-Xin Ma
Approved by the Committee, November, 2014
Committee Cnair Professor Gunter Luepke, Applied Science
The College of William and Mary
Professor Michael Kelley, Applied Science The College of William and Mary
Professor Mark Hinders, Applied Science The College of William and Mary
Professor Mumtaz Qazilbash, Physics The College of William and Mary
ABSTRACT
This thesis presents optical characterization of the static and dynamic magnetic interactions in ferromagnetic and multiferroic heterostructures with time-resolved and interface-specific optical techniques. The focus of the thesis is on elucidating the underlying physics of key physical parameters and novel approaches, crucial to the performance of magnetic recording and spintronic devices.
First, time-resolved magneto-optical Kerr effect (TRMOKE) is applied to investigate the spin dynamics in L1o ordered FePt thin films, where perpendicular magnetic anisotropy Ku and intrinsic Gilbert damping ar0 are determined. Furthermore, the quadratic dependence of Ku and a0 on spin- orbit coupling strength f is demonstrated, where f is continuously controlled through chemical substitution of Pt with Pd element. In addition, a linear correlation between cro and electron scattering rate 1 /re is experimentally observed through modulating the anti-site disorder c in the L1 0 ordered structure. The results elucidate the basic physics of magnetic anisotropy and Gilbert damping, and facilitate the design and fabrication of new magnetic alloys with large perpendicular magnetic anisotropy and tailored damping properties.
Second, ultrafast excitation of coherent spin precession is demonstrated in Fe/CoO heterostructures and Lao.6 7Cao.3 3Mn0 3 thin films using TRMOKE technique. In the Fe/CoO thin films, Instant non-thermal ferromagnet (FM) - antiferromagnet (AFM) exchange torque on Fe magnetization through ultrafast photo-excited charge transfer possesses in the CoO layer is experimentally demonstrated at room temperature. The efficiency of spin precession excitation is significantly higher and the recovery is notably faster than the demagnetization procedure. In the Lao.67Cao.33Mn03 thin films, pronounced spin precessions are observed in a geometry with negligible canting of the magnetization, indicating that the transient exchange field is generated by the emergent AFM interactions due to charge transfer and modification of the kinetic energy of eg electrons under optical excitation. The results will help promoting the development of novel device concepts for ultrafast spin manipulation.
Last, the interfacial spin state of the multiferroic heterostructure PbZro.5 2Tio.4 8O3/Lao.6 7Sro.3 3MnO3 and its dependence on ferroelectric polarization is investigated with interface specific magnetization induced second harmonic generation (MSHG). The spin alignment of Mn ions in the first unit cell layer at the heterointerface can be tuned from FM to AFM exchange coupled, while the bulk magnetization remains unchanged as probed with MOKE. The discovery provides new insights into the basic physics of interfacial magneto-electric (ME) coupling.
TABLE OF CONTENTS
Acknowledgements iii
Dedication iv
List of Tables v
List of Figures vi
Chapter 1. Introduction
1.1 Magnetic recording and spintronics 1
1.2 Magnetization switching and dynamics 3
1.3 Interface magnetization 6
1.4 Scope of dissertation 9
Chapter 2. Experimental techniques
2.1 Magneto-optical Kerr effect 12
2.2 Time-resolved Magneto-optical Kerr effect 17
2.3 Magnetization induced second harmonicgeneration 26
Chapter 3. Time resolved studies of spin dynamics in FePt thin flims
3.1 Introduction 34
3.2 Samples and characterizations 36
3.3 Magnetic anisotropy and damping inFePtxPd(i-x) thin films 43
3.4 Magnetic anisotropy and damping in FePtthinfilms with different chemical order 55
Chapter 4. Time resolved spin precession in Fe/CoO heterostructure
4.1 Introduction 67
i
4.2 Samples and characterizations 69
4.3 Magnetic anisotropy 72
4.4 Ultrafast spin exchange torque via photoexcited charge transfer processes 82
Chapter 5. Spin precession excitation in manganite thin films
5.1 Introduction 100
5.2 Samples and experiments 102
5.3 Coherent Spin Precession via Photoinduced Anti-ferromagnetic Interactions in Lao.6 7Cao.3 3Mn0 3 103
Chapter 6 . Magneto-electric coupling at PbZr0 5 2Ti 0 .4 8 O3 /Lao 6 7Sr0.3 3MnO3 interface
6.1 Introduction 117
6.2 Samples and experiments 119
6.3 Charge control of antiferromagnetism atPbZr0 52Ti 0 4 8 O3/ La0.67Sro.33Mn03 interface 124
Chapter 7. Conclusions 136
Bibliography 139
Vita 156
ii
ACKNOWLEDGEMENTS
I am truly grateful to everyone who has directly or indirectly supported and helped me during the completion of this dissertation. First and foremost, I would like to thank my advisor, Professor Gunter Luepke, for his guidance, encouragement and mentoring throughout my doctoral study at the College of William and Mary. He has always provided enthusiasm and perspective for the research process. This dissertation would not have been possible without his support and guidance.
I am particularly grateful to Professor Haibin Zhao for guiding my experiment in the lab and fruitful scientific discussions. I truly appreciate his sharing of academic experience with me as a friend.
I owe special thank to Professor Shiming Zhou and his group members - Pan He and Li Ma at Tongji University in Shanghai, China. I am indebted to them for providing us state-of-art FePt thin films, structure characterizations and inspired visions.
I truly appreciate Professor Yizheng Wu and his group members- Qian Li and Jie Zhu at Fudan University in Shanghai for their support and providing us high-quality Fe/CoO thin films.
I am similarly grateful to Professor Ram Katiyar and Dr. Ashok Kumar for providing us high-quality multiferroic heterostructures and characterization results. I also appreciate Professor James Scott for helpful suggestions.
I am indebted to Professor Qi Li from Penn State University, for her generosity in providing us high-quality manganite films and sharing academic experience.
My sincere gratefulness goes to my lab mates: Fan Fang, Wei Zheng, Haowei Zhai. I appreciate all their friendships and their collective encouragement to finish this dissertation. I have enjoyed every moment that we have worked together.
I would also like to express my appreciation to my doctoral dissertation committee members: Professor Michael Kelley and Professor Mark Hinders of Applied Science, Professor Mumtaz Qazilbash of Physics for their comments and suggestions during the course of completing this dissertation.
Finally, I wish to express my sincere gratitude to my parents - Renmi Ma and Ling Huang, and my wife - Yufei Ding. Their constant love, encouragement and support have always been the source of my strength and motivation. Anything I have been able to accomplish is a tribute to them.
LIST OF TABLES
2.1 Fresnel coefficients, where = cosO, 6 the light incidentangle, a2 = J l - sin20 /n 2, QP,QT,QLthe polar, transverse,and longitudinal projections of the Voigt vector Q, respectively.
2.2 Nonzero tensor elements of the nonlinear susceptibility tensor * (2) for the (0 0 1) cubic surface, the magnetization parallel to the x-axis (longitudinal configuration), y-axis (transversal configuration) and parallel to the z-axis (polar configuration). * (+) and denote the even and odd elements with magnetization. Other non-zero elements with mixed polarization of s and p are not included. For simplicity of notation the tensor components are indicated by their indices only
3.1 Landau g factors and Ku derived from TRMOKEmeasurements for different doping level x at 20K and 200K.
LIST OF FIGURES
1.1 Schematics of magnetic recording arrays with traditional writing mechanism.
1.2 Schematic of precessing magnetization.
1.3 Relative orientations of the atomic moments in the AFM and FM are shown schematically, illustrating the lateral offset in the magnetization curve.
2.1 Schematics of longitudinal, transverse and polar MOKE probing geometry, where POS is the plane of sample (x-y plane), and POI is the plane of incident (y-z plane).
2.2 Geometry of TRMOKE measurements.
2.3 Stages of the excitation process: (a) M aligns along the equilibrium direction H eff \ (b) new equilibrium directionH'eff when pump pulse “heat up” the sample; (c) Mprecesses around H'eff (d) heat diffused and Mprocesses around H eff.
2.4 TRMOKE result from L10 FePt thin film with growth temperature 550 °C at magnetic field 3.5T.
2.5 Schematic of TRMOKE measurement geometry for L1 o FePt thin films.
2.6 Coordinate system used for calculating the magnetization precession frequency in the L10 FePt thin films.
2.7 Schematic of TRMOKE experimental set-up.
2.8 Schematic of MSHG and MOKE measurements on PZT/LSMO thin film.
2.9 MSHG hysteresis loop from PZT/LSMO thin films indicating l(-M) and l(+M) used to determine the magnetic contrast A from Eq. (2.17).
2.10 The schematic set-up of MSHG measurement.
3.1 Schematics of atomic structures for the two series ofFePt thin films.
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3.2 Structure characterization results of FePtxPd(i.X) thin films(a) and FePt thin films with different Tg (b) by X-ray Diffraction measurements.
3.3 Anti-site disorder percentage c in Z_10 ordered FePt thin films as a function of growth temperature Tg.
3.4 Static magnetic hysteresis loops measured by vibrating sample magnetometer (VSM) (a)-(c) for FePtxPd(i-x) thin films. The black (red) curves represent out-of-plane (inplane) results. VSM results for FePt thin films with different Tg (d).
3.5 TRMOKE data with different doping level of FePtxPd(l-x) under magnetic field H=5T. The solid lines are fitted curves.
3.6 TRMOKE results of FePt0.82Pd0.18 (a) and FePt0.5Pd0.5 (b) under different magnetic fields H. The solid lines are fitted curves.
3.7 The dependence of precession frequency (a) and lifetime(b) on external magnetic field H for typical samples. The solid lines refer to fitted results.
3.8 PMA (a) and ao (b) as a function of doping level x of FePtxPd(i-x)
3.9 Spin-orbit coupling strength f (a), bandwidth W (b), spin moment ms on Fe or Pt/Pd site (c) and measured c/a ratio (d) as a function of doping level x of FePtxPd(i.X).
3.10 Measured and calculated ao (a) and PMA (b) as a function of SOC strength. The solid line is a quadratic fit obtained from first principle calculations.
3.11 TRMOKE data from 22-nm thick FePt films, grown at different Tg with applied magnetic field H=6.5T (a), and with Tg =680 °C measured at different H (b). The solid lines refer to fitted results.
3.12 The dependence of spin precession frequency f (c) and effective Gilbert damping aeff (b) on H obtained from Eqs.3.3 and 3.5 for FePt thin films with different x. The solid lines are fitted results.
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3.13 PMA Ku (a) and Landau g factor (b) as a function of antisite defect concentration c for FePt films with thicknesses of 17 nm and 22 nm.
3.14 ao (a), and ao/g (b) as a function of anti-site defect concentration c for FePt films with thicknesses of 17 nm and 22 nm. The solid line refers to least square linear fit.
3.15 The dependence of f and 1/r on temperature T with applied field H =5T (a), and as a function of H at 20 K (b). The solid lines refer to fitted results.
4.1 Longitudinal hysteresis loops for magnetic field along uniaxial easy axis [100] and hard axis [010] at 82 K and 330 K.
4.2 Temperature dependence of coercivity (Hc) and exchange biasing field (He) obtained from easy axis loops.
4.3 TRMOKE results of Fe/CoO thin films with different magnetic field H along the Fe [110] direction (in-plane hard axis) at RT.
4.4 Spin precession frequency f as a function of magnetic field H.
4.5 Temperature dependent TRMOKE results with magnetic field H= 3 kOe.
4.6 Temperature dependent TRMOKE results with magnetic field H= 1 kOe.
4.7 Spin precession frequency fas a function of temperature.
4.8 (p as a function of temperature (green=3 kOe, blue=1 kOe).
4.9 9 as a function of temperature (green=3 kOe, blue=1 kOe).
4.10 Uniaxial magnetic anisotropy field KJMs as a function of temperature.
4.11 Two pump strategies to optically excite the spin precession, where the black arrows represent the spins of magnetic atoms and ions. The optical charge transfer transition in CoO is depicted.
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4.12
4.13
4.14
4.15
4.16
4.17
4.18
4.19
5.1
5.2
5.3
TRMOKE results from Fe/MgO (black squares) and Fe/CoO (red circles) with pump pulse intensity 3.1 mJ/cm2, and Fe/CoO (blue triangles) with pump intensity 0.16 mJ/cm2.
Measured and simulated spin precession amplitude A as a function of T.
The spin precession amplitude A is plotted as a function of H at different T, where solid lines represent simulated results.
Measured Kerr signal change with 0.8 ps step size (black squares), simulated polar magnetization component (red curve) and assumed modulation of Ku (blue curves) as a function of t.
Measured Kerr signal change with 4 ps step size (black squares), simulated polar magnetization component (red curve) and assumed modulation of Ku (blue curves) as a function of t.
Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with 77 = 2 ps.
Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with 7> = 15 ps.
Measured Kerr signal change (black squares) and simulated polar magnetization component (red curve) with rr = 1 0 0 ps.
Transient reflectivity AR (a), and transient Kerr rotation A0 of LCMO film with magnetic field H (0.5 T) applied perpendicular to the sample plane (b), with H applied along the in-plane easy axis (c), and with zero field (d) at 20 K; Fourier transform (e) of Fig. 5.1(d).
Transient Kerr rotations of Fe film with magnetic field applied along the in-plane hard axis (a) and easy axis (b) at RT.
Three-temperature model in ferromagnetic metals and half-metals.
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5.4 Amplitude of precession of 60 nm and 100-nm LCMO films with field applied at 45 degree to the easy axis and along the easy axis at 20 K (a), and Fe film with field applied at 55 degree to the easy axis and of 10 degree to the easy axis at RT (b).
5.5 Transient exchange field (a) and transient anisotropy field (b) in a state of canted magnetization (magnetic field at 5 degree to the easy axis) of 60-nm thick LCMO film. The dashed line is a guide to the eye of the data.
5.6 Transient Kerr rotations of LCMO film with applied field H=0.5T at various temperatures. The inset shows the uniform precession frequency as a function of the temperature.
6.1 Schematic of MSHG and MOKE measurements (a), and depiction of PZT/LSMO interface on atomic scale (b). MSHG probing plane is marked with purple color.
6.2 XRD results of PZT/LSMO heterostructure (a), RHEED patterns of LAO substrate surface (b), LSMO surface before deposition of PZT layer on top (c) and PZT surface of PZT/LSMO/LAO heterostructure (d).
6.3 Interface (a) and Bulk (b) magnetic hysteresis loops from PZT/LSMO heterostructure for different external gate voltages Ug measured with MSHG and MOKE techniques at 78 K, respectively.
6.4 MSHG hysteresis loop with Ug = +10V indicating l(-M) and l(+M) used to determine the magnetic contrast A from Eq. (6.1) (a), and MSHG magnetic contrast A as a function of Ug (b). Increasing (decreasing) gate voltages, Ug up (Ug down), are labeled in red (black). The direction of positive Ug is defined in the inset.
6.5 Ferroelectric polarization (P) in PZT as a function of Ug (a). The direction of positive P is defined as pointing towards LSMO thin film. MSHG magnetic contrast A as a function of P of PZT (b). The dashed line in Fig. 6.5(b) is the least square fit for the data points.
6.6 Strain state of LSMO interface as a function of gate voltage (a). MSHG magnetic contrast A as a function of strain (5) of PZT (b).
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6.7 Schematic model of ions displacement and screen charge accumulation under opposite polarization state.
6.8 P in S out MSHG results under different polarization states.
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Chapter 1
Introduction
1.1 Magnetic recording and spintronics
For the past few decades, magnetic recording and spintronics have
been the major technologies for information storage and processing [1-2]. The
magnetic devices have been widely applied in the hard disk read heads,
memories and sensors, where the binary information is represented by the
magnetization of magnetic elements, instead of electrons or holes in
traditional semiconductor electronics [3]. Nowadays, the large amount of
digital information creates ever-increasing demands for much higher storage
densities, faster and more energy-efficient processing operations. Therefore,
it attracts abundant research interest to investigate the key physical
parameters and mechanisms for the development of device performance.
The storage density is closely related to the magnetic anisotropy of the
magnetic element [4-6]. Each piece of magnetic recording media carries a bit
of information through controlled magnetization directions, while the element
size is limited by the thermal fluctuations of the magnetization directions [7, 8].
Magnetic anisotropy provides the energy barrier between these states,
allowing them to withstand the loss of stored information due to thermal
fluctuations. Therefore, the search for material structures with larger magnetic
anisotropies promotes the increase of storage densities, as shown in Fig 1.1.
l
In order to improve the speed and energy consumption of the
processing operations, novel approaches to effectively manipulate the
magnetization in the magnetic devices are investigated. For example,
optically generated effective magnetic field pulses and the spin-transfer
torque (STT) effect from spin-polarized carrier transport are recently applied
to switch the magnetization in spin valves and magnetic tunneling junctions
(MTJs) [9-12]. These techniques can be very strong, interact locally [13], and
thus are favorable for the magnetization manipulation in high-density
magnetic devices. With respect to energy consumption, it is more energy-
favorable to generate electric fields instead of currents in highly integrated
devices. As a result, the electric control of magnetization offers new pathways
to improve device performance. [14] Multiferroic materials and
heterostructures are of particular interest for using electric fields to control the
magnetization via the ME coupling effect [15].
Figure 1.1 Schematics of magnetic recording arrays with traditional writing
mechanism [16].
*Ring* writing element
Longitudal Recording (standard)
] Recording layer
'Monopole’ writing element
Perpendicular Recording
La Recording Layer | Additional Layer
2
1.2 Magnetization switching and dynamics
1.2.1 Magnetization switching process
In the demand for faster operation, the dynamic properties of
magnetism in magnetic devices and controlled magnetization switching
become increasingly important. The magnetization switching process is well
modeled with the Landau-Lifshitz-Gilbert (LLG) equation [17],
d M m m > • M C^M / jt jk \— = - YM x H eff + a - x — (1.1)d t 1 M d t ' ’
where the first term on the right indicates that the magnetization M does not
go directly to the desired alignment along H eff, but instead, circles around
H ef f ■ The second term describes the damping of the circling precession
towards H eff with the Gilbert damping parameter a, where larger a
represents faster relaxation and dissipation of magnetic energy into the
thermal bath. Therefore, the speed and damping of coherent spin precession
play important roles in the operation speed of magnetic devices.r
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Figure 1.2 Schematic of precessing magnetization.
3
1.2.2 Damping mechanisms
Well-controlled magnetic relaxation is required for designing magnetic
devices with promising performance. For most magnetic transition metals, a is
small; so that M precesses back and forth several times after switching or
small perturbation, which takes several nanoseconds to relax the precession.
This actually puts a bottleneck on the switching speed. Therefore in many
applications requiring high speed or high data rates, the stronger damping
behavior is desirable. On the other hand, larger cr indicates stronger coupling
between the magnetic system and the thermal bath. This would allow the
thermal bath to drive fluctuations of the magnetization, and cause
magnetization noise in small magnetic sensors. For these applications, the
noise limit could be improved with reduced damping. Thus in all, some
sophisticated methods to control the magnetic relaxation behavior are desired.
The Gilbert damping parameter a contains both extrinsic and intrinsic
components. The extrinsic Gilbert damping is due to nonlocal spin relaxation,
such as spin pumping and two-magnon scattering. This extrinsic component
can be tuned by artificial substrates, specially designed buffer and coverage
layers. On the other hand, the intrinsic Gilbert damping parameter a0 is due to
a local magnetic energy transfer effect and describes the energy flow rate
from spin to electronic orbital and phonon degrees of freedom through
electron scattering.
4
1.2.3 Excitation of Magnetization switching
The conventional method to switch the magnetization alignment is
magnetizing the storage cell with an external magnetic field. As mentioned
above, the increasing storage density requires larger magnetic anisotropy,
which provides a significant energy barrier for switching the magnetization
between the spin-up and spin-down states. The switching actually requires
the application of stronger magnetic field, leading to larger energy
consumption. Recently, a breakthrough of this limitation is the utilization of
optical demagnetization or spin-polarized current to overcome the large
magnetic energy barrier. For instance, effective magnetic field pulses can be
generated by the optical modulation of magnetic anisotropies with laser
pulses, which trigger the magnetization switching. Another example is the
spin-transfer torque (STT) effect generated by optically or electrically
triggered spin-polarized carrier transport. The STT effect is due to the
interaction between itinerant electrons in FM thin films and the magnetization
[13], and can be used to switch the magnetization.
The critical density of the spin polarized current for STT switching is
proportional to aKuMs [9-12], where a is the Gilbert damping parameter, Ku
the magnetic anisotropy and Ms the saturated magnetic moment. Considering
the storage density, the speed and energy consumption of magnetization
switching in STT devices, a smaller or along with large value of Ku is essential
for promising performance.
5
1.3 Interface magnetization
1.3.1 Magneto-electric coupling
The switching of magnetic state is accomplished by applying an
electric current, either for the generation of a magnetic field or for the STT
effect. The large current densities required also lead to significant energy loss
from heating. In contrast, it is more convenient and energy-favorable to
generate an electric field in high-integrated devices. Thus in the push for low-
energy-consumption logic devices, strong ME coupling seems to offer a new
mechanism to break through those limitations and improve the device
performance; as such, it has recently drawn increasing research interest.
In the search for stronger ME coupling, multiferroic structures are
promising candidates, where the materials exhibit both FM and ferroelectric
(FE) order. These materials are usually transition metal oxides, where the
competition between different interactions in such strong-correlated electron
systems results in a complex phase diagram, and provides a platform to
manipulate the coupling between different order parameters. For example in
some manganites [18], there are some phases with the coexistence of FM
and FE order. However, such phases usually exist at low temperature, and
the ME coupling is fairly weak, which prevents them from real-life applications.
To overcome these disadvantages, an alternative but challenging approach is
the ME coupling across the interfaces of multiferroic heterostructures, which
contain room-temperature FE and FM order. Moreover, the electric control of
6
interface magnetization can be effectively incorporated into devices where the
performance depends crucially on the interface magnetization, such as MTJs.
FM
A FM
(ii)
(iii)
(Hi) (iv)
(iv )
-> H
Figure 1.3 Relative orientations of the atomic moments in the AFM and FM are
shown schematically, illustrating the lateral offset in the magnetization curve.
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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 7Cao3 3 Mn0 3, due to its complex
magnetic phase diagram. For the first time, pronounced coherent spin
precession is observed in a geometry with negligible canting of the
magnetization in ferromagnetic Lao6 7Ca0.3 3Mn0 3 thin films using TRMOKE
technique. This indicates that the transient exchange field is generated by the
emergent AFM interactions due to charge transfer and modification of the
kinetic energy of eg electrons under optical excitation.
Multiferroic complex oxide heteroiunction - the interfacial spin state of
the multiferroic heterostructure PbZr0.52Ti0.48O3/La067Sr0.33MnO3 and its
dependence on ferroelectric polarization is investigated with interface-specific
MSHG technique. The spin alignment of Mn ions in the first unit cell layer at
the heterointerface can be tuned from FM to AFM exchange coupled, while
the bulk magnetization remains unchanged probed with MOKE.
10
Chapter 2
Experimental techniques
This chapter presents three different optical characterization
techniques which have been employed to study magnetic and electronic
properties: MOKE, TRMOKE and MSHG.
Section 2.1 describes the MOKE technique and the optical geometries
for measuring magnetization vector components. The principles and
experimental setup of TRMOKE technique - an all-optical technique to study
spin dynamics - is described in section 2.2. The magnetization precession
and its mathematical derivation are also discussed in this section. Section 2.3
presents an interface-specific magnetization probe technique - MSHG,
including a detailed analysis of nonlinear susceptibilities in PbZro 5 2Tio 4 8 O3
/La0 6 7 Sro.3 3Mn0 3 and corresponding experimental consideration for probing
interface magnetization.
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2.1 Magneto-optical Kerr effect
In general, the polarization state of light reflected from a magnetized
material will change depending on the state of magnetization, resulting in a
modified ellipticity [20]. If the incident light is linearly polarized, this process is
known as MOKE, which was first discovered by John Kerr in 1877 [21]. The
origin of this effect is presently described in the context of either microscopic
quantum theory [22] or macroscopic dielectric theory [23],
Microscopically, this effect originates from spin-orbit coupling between
light and electron spin. A potential gradient \7V is generated by light, which
couples local spins,5, with momentum,P, adding an interaction term (W x
P) • S to the Hamiltonian [22], In the macroscopic picture, such interaction
gives rise to off-diagonal components of the dielectric tensor F, which are
determined by the state of magnetization in the magnetic medium:
1 iQz - iQ y - iQ z 1 iQxiQy iQx 1
(2 .1)
where e is the average dielectric constant, and Q = ( Qx, Qy, Qz) is the Voigt
vector, which is proportional to the magnetization [24].
From Maxwell’s equations, the normal modes of light propagation can
be categorized into left-circular polarized light with index of refraction nL =
n ( l - i <?•£), and right-circular polarized light with refraction index nR =
n( 1 +^Q • k), where n = V i is the average refraction index and k is the unit
12
vector of the light propagation direction. A linear polarized light can be
considered as a superposition of left-circular and right-circular polarized light
with equivalent amplitude. The difference between the real parts of nL and nR
in the medium will lead to a phase shift between the two modes and result in
a polarization rotation of the light propagating through the medium.
Meanwhile, the difference between imaginary part of nL and nR will cause the
different absorption rates for the two normal modes, affecting the ellipticity of
the light.
In the Kerr effect, a convenient decomposition scheme is that the
polarization is a combination of p- and s-polarized light, which define the
polarization of light parallel and perpendicular to the incident plane,
respectively. Therefore, MOKE can be represented by:
where Eg1 and £ jnare the p- and s-polarized components of incoming light,
Eput and E$ut are corresponding components for outgoing light, and the
matrix r is the scattering matrix with Fresnel coefficients.
MOKE can be categorized into longitudinal, transverse, and polar
MOKE, according to the geometry of magnetization (M) with respect to the
incident plane of light and film plane (Fig. 2.1): in the longitudinal geometry
the magnetization component lies in both the light incident plane and film
plane; in the transverse geometry the magnetization component is in the film
(2.2)
13
plane but perpendicular to the light incident plane; for the polar geometry, the
magnetization component is perpendicular to the film plane.
For different MOKE geometries, the Fresnel coefficients are distinct, as
summarized in Table. 2. 1. As a result, the longitudinal and polar components
of M change the polarization state of the light upon reflection. In contrast, the
transverse component of M only changes the intensity of the reflected beam,
and no ellipticity occurs. Therefore, different detection schemes are required
to measure different magnetization components. For example, for p-polarized
incoming light the change of s-polarization is sensitive to the longitudinal and
polar components of M , while one has to probe the change of p- polarization
to measure the transverse component of Af.
For the MOKE measurements, a beam of light first passes through a
linear polarizer and is then incident on the sample; after reflection from
sample, the beam passes through an analyzing polarizer (analyzer) and is
measured by a photodetector. In all my measurements, the cross-polarized
combination of polarizers is used to detect the longitudinal and polar
components of M. In other words, the analyzer angle is set close to 90° with
respect to the incident beam polarization.
14
polar
Figure 2.1 Schematics of longitudinal, transverse and polar MOKE probing
geometry, where POS is the plane of sample (x-y plane), and POI is the plane of
incidence (y-z plane).
15
Table 2.1: Fresnel coefficients, where ax = cos0, 6 is the light incident angle,
a 2 = yjl - sin26 / n 2, QP, QT, QLare the polar, transverse, and longitudinal projections
of the Voigt vector Q, respectively.
Component vv YpS — r sp ss
j-i,
Longitudinal
* ie w + * * i. ' '* wt 1 ,A' ■ . * a1-n a 2
n a 1- a 2n a x+ a 2
•. jf(naa+ a z) (n a 2+ a i) a t+ n a 2
Qf,na1tan02 «1-n a2i(na1+a2)(na2+a1) ax+na2
-WW...
i ;t ȣ} pry* ' ^
l^ C ^ /o ^ a i—thQ^itanBz ■ *
T ~~ . /.'.w*? ' <• I/ “ • I ;X ; i l
ax-n a 2ax+na2
16
2.2 Time- resolved Magneto-optical Kerr effect
TRMOKE is an all-optical method to investigate the spin dynamics in
the time domain [25], which utilizes the pump-probe technique. The pump
beam is used to excite the magnetization precession, while the probe beam
measures the decay of excited magnetization state with MOKE technique. By
controlling the delay between pump and probe laser pulses, a real-time plot of
magnetization precession is obtained.
_ PumpProbe TRMOKE
Substrate
Figure 2.2 Geometry of TRM OKE measurements.
17
2.2.1 Excitation of magnetization precession
Figure 2.3 depicts the stages of the excitation process by pump laser
pulses. Before the pump pulse reaches the sample t< to , the magnetization Af
lies along the equilibrium direction H eff , determined by the magnetic
anisotropy field, demagnetization field, exchange coupling, and external
magnetic field H. Here we assume the easy axis of magnetic anisotropy is
perpendicular to the film plane. When the pump pulse impinges on the
sample at t= to , the magnetic state of the sample is modified, which can affect
the magnetic anisotropy, exchange coupling, and/or the demagnetization field.
As a result, the equilibrium direction of M a t this stage is changed
which induces a torque on M to align it along the new direction. Within a short
time region fo^fo+IO ps, M precesses around the new direction. After the
original magnetic state of the sample is restored, t> fo+ 1 0 ps, the precession
continues around the equilibrium direction due to the initial displacement of M.
Figure 2.4 displays the TRMOKE result from L1o FePt thin films. The
excitation of magnetization precession is as depicted in Fig. 2.3. In the
TRMOKE measurement, we utilize the MOKE geometry sensitive to polar
component (p-in, s-out). Therefore, Figure 2.4 can also be interpreted as a
real- time plot of Mz as a function of delay time. The t m is the time for
demagnetization process, corresponding to Fig. 2.3 (b). During this process,
the electron temperature Te is first raised by photon-induced excitations (spin-
conserved), and then relaxed through the phonon assisted spin-flip scattering.
18
Meanwhile, the spin temperature Ts increases, resulting in a reduction of
magnetization. The te is the time for the recovery process, depicted in Fig.
2.3 (c). Afterwards, the magnetization starts to precess around the equilibrium
direction as depicted in Fig. 2.3 (d).
t> to + 1 0 psto < t< to + 1 0 ps
Figure 2.3 Stages of the excitation process: (a) M aligns along the equilibrium
direction H ef f \ (b) new direction H'eff when pump pulse “heat up” the sample; (c) M
precesses around H'e ff (d) heat diffused and M precesses around H e ff .
19
O)
0 20 40D elay tim e (ps)
Figure 2 .4 TRM OKE result from /_10 FePt thin film with growth temperature 550 °C at
magnetic field 3.5T.
2.2.2 Precession frequency and effective Gilbert damping
The precession motion of Af back to the equilibrium direction can be
phenomenologically modeled with LLG equation [17]:
where y = ggB/h is the gyromagnetic ratio, and a is the dimensionless Gilbert
damping constant. The effective field H eff is the sum of external magnetic
20
field, demagnetization field, anisotropy field, dipolar field, and exchange field.
The first term on the right describes the torque on the magnetic moment
produced b y H eff, which drives a precession motion around the direction of
the equilibrium direction H e/ / . The second term on the right describes the
damping of precession due to the exchange of energy and angular
momentum with environment, until the magnetization finally aligns parallel to
H eff . Figure 2.5 depicts an example of magnetization precession.
Figure 2 .5 Schematic of TRM OKE measurement geometry for L10 FePt thin films.
For small angle precession, Eq. (2.3) can be solved analytically with
the following steps: ( 1 ) finding the equilibrium orientation of magnetization
determined by H eff\ (2 ) deriving the equations of motion by describing the
torque on spins by the effective field; (3) solving the linearized equations of
motion, to obtain the frequencies and Gilbert damping of the spin wave
modes.
21
For example in L1o ordered FePt material system, the magnetic free
energy F is the sum of Zeeman energy, demagnetization and anisotropies:
F = —HMS cos(0 — (p) + 2ttM^ c o s 2 0 + Kusin20 (2.4)
where H is the magnitude of applied magnetic field, Ms is the saturated
magnetization, B and q> are the angles between M and [0 0 1 ] direction, and H
and [0 0 1 ] direction, respectively. Ku is the perpendicular magnetic anisotropy.
For convenience, the anisotropy field Ha = 2Ku/M s and demagnetization field
Hd = 4tcMs . The equilibrium direction for M can be derived by taking the
derivative of F over 0:
n _ s in (20)Ha- H d 2*s in ( (p -6 ) ' ' ’
i L
Z ’ //[001]z / / / I I
X ®
Figure 2 .6 Coordinate system used for calculating the magnetization precession
frequency in the L10 FePt thin films.
22
The magnetization precession motion can be written as:
d m x •* , d m y d m z—~ x + ——y + —d t d t d t
-»X y-*z
y mx my Ms + mzHx Hy Hz
(2.6)
where Hx,Hy,Hz are the x, y, z components of H efr, and can be derived by:
dF
d M y
Hy = Hs\n(0 - qi) (2.7)
d F
d M ,+ Heos(0 - <p)
Therefore, the torque equations become:
1 drriy
y d t= — [H cos(0 — <p) + {Ha — Hd) * cos2B]my
^ = [H cos(0 -<p) + (Ha - Hd) * cos(29)]mx (2 .8)
For the uniform precession, the my and mz can be written as:
my = m yQe l2TCft+<py
m z = m z0e i 2 n f t+ < p z (2.9)
The precession frequency f of spin wave can be derived by substituting Eq.
(2.9) into Eq. (2.8):
2tt/ = yV[W cos(0 - <p) + (Ha - //d) * cos(28)][H cos(0 - <p) + {Ha - Wd) * cos20]
(2 .10)
23
In terms of Gilbert damping parameter a, one has to incorporate the
damping term in Eq. (2.6), and the uniform magnetization precession can be
described by an oscillating term times the decay term where
a y [ ( H cos(0 - <p) + ( tfa - tfd) * cos20 ) + ( / / cos(0 - <p) + ( t fa - tfd) * cos(20))]
T “ 2(1 + a 2)
(2 .11)
is the lifetime of magnetization precession. Therefore, given the precession
frequency f and lifetime r as a function of magnetic field H, we can derive
some physical parameters of the magnetic material, such as magnetic
anisotropy and effective Gilbert damping parameter aeff.
2.2.3 Experimental set-up
Figure 2.7 displays the schematic TRMOKE setup used in this
dissertation. TRMOKE measurements are performed in a pump-probe setup
using pulsed Ti.sapphire laser (RegA-9000 amplifier laser system) with a
pulse duration of 200 fs and a repetition rate of 250 kHz. The wavelength of
pump (probe) pulses is 400 nm (800 nm). A modulated pump pulse beam
with a fluence of 0.16 mJ/cm2 is focused to a spot ~ 1 mm in diameter on the
sample to excite the magnetization precession, and the transient Kerr signal
is detected by a probe pulse beam which is time-delayed with respect to the
pump. The focus area of probe beam is with a diameter of 0.7 mm, which is
24
smaller than that of pump beam; so that the intensity ratio of the pump to
probe pulses is set to be about 6:1. A variable magnetic field H is applied
using a superconducting magnet produced by Oxford Instruments. The probe
signal is send to a Lock-in amplifier produced by Stanford Instruments.
Polarizer (P)
Analyzer (S)
Electromagnet
Figure 2.7 Schematic of TRM OKE experimental set-up.
25
2.3 Magnetization-induced second harmonic generation
MSHG is a nonlinear optical technique that probes the interface or
surface magnetization [26], and sometimes it is called nonlinear MOKE. The
MSHG signal can be generated in magnetic materials in which both space-
inversion and time-reversal symmetry are simultaneously broken. The broken
time-reversal symmetry results from the presence of magnetization, while the
broken space-inversion symmetry occurs only at the surface or interface of
centrosymmetric materials. Figure 2.8 depicts the different probing regions of
MOKE and MSHG in La0.6 7 Sr0.3 3 MnO3 magnetic thin film capped with
PbZr0 5 2Tb 4 8O 3 non-magnetic layer. MOKE probes the bulk magnetization,
while MSHG probes the interface magnetization.
MSHG MOKE
Interface
LSMO
Figure 2.8 Schematic of MSHG and MOKE measurements on PZT/LSM O thin film
26
2.3.1 Second harmonic generation (SHG)
SHG signal arises from the nonlinear polarization P(2oS) induced by an
incident laser field E(co), where co is the angular frequency of light. This
polarization can be written as an expansion in E(co):
P(2o>) = x {2)E{ai)E(a>) + x (Q}E(o))VE(a)) + - (2.12)
The lowest-order term in Eq. (2.12) describes an electric dipole source, with
* (2) as a third-rank susceptibility tensor. Symmetry considerations show that
this contribution is zero in a centro-symmetric medium, thus limiting electric
dipole radiation to the surface or interface where space-inversion symmetry is
broken. The second term in Eq. (2.12) describes the bulk SHG in terms of the
much smaller electric quadrupole-like contributions, where * (<2) is fourth-rank
susceptibility tensor. However, because of the large volume difference
between interface and bulk this does not necessarily mean that the total bulk
second-harmonic signal is negligible. Moreover, electric field induced SHG is
also generated in the bulk of centrosymmetric materials, and this effect can
be significant since the quasi-static electric field can reach very high values.
[27]
2.3.2 Nonlinear magnetic susceptibility
The presence of magnetization will not break the space-inversion
symmetry of the bulk, but can lower the surface and interface symmetry
modifying the nonlinear susceptibility. As a result, the surface or interface
27
second harmonic polarization can be simplified to one third rank tensor with
different components:
Pt( 2a>) = <JCijkw ( ± M ) + Xijk(-H±M))Ej(a>)Ek(a>) (2.13)
where i, j, k=x, y or z indicate the directions of longitudinal, transverse and
polar components, * ijk(+) and represent the tensor components even
and odd in M, respectively. The odd component ^jjk(_) changes sign upon
magnetization reversal and therefore give rise to the magnetic asymmetry in
the MSHG response.
The symmetry of the nonlinear susceptibility is determined by the
symmetry of the particular surface under consideration. Some components of
X are non-zero, if they satisfy the same space symmetry. Therefore, the non
zero elements of x can be obtained from invariance of x under symmetry
operations:
Xi \k = Tii'VTkk'^i'j'k' (2-14)
where T is the transformation matrix for each symmetry operation, and
summation over repeated indices is implied. Take the magnetized (001)
surface with cubic symmetry for example: With M aligned along [100] (x)
direction, the two independent symmetry operations are reflection about the
x-z mirror plane, (x ->x, y ->-y, and Af-> -M) and reflection about y-z mirror
plane, (x ->-x, y ->y, and M). By utilizing Eq. (2.14), one can obtain the
non-vanishing elements. Xxxz> Xxzx> Xyyz> Xyzy> Xzxx> X z y y 3nd Xzzz< which are
28
even terms, x Xxy> Xxyx> Xyxx> Xyyy> Xyzz> X z y z X z zy odd terms. Table
2.2 summarizes the non-zero terms in different cases.
Table 2.2 Nonzero tensor elements of the nonlinear susceptibility tensor * (2) for the
(001) cubic surface, the magnetization parallel to the x-axis (longitudinal
configuration), y-axis (transversal configuration) and parallel to the z-axis (polar
configuration). and denote the even and odd elements with magnetization.
Other non-zero elements with mixed polarization of s and p are not included. For
simplicity of notation the tensor components are indicated by their indices only.
InputPolarization
OutputPolarization
Longitudinal
J lf//x
yzz, yxx P s
zzz, xzx xxz, zxx
P P
zyy s Pyyy s s
Transverse
H f//y
P sxxz, xzx zxx, zzz
zxz, zzx XXX, xzz
P p
zyy xyy s ps s
Polar
M / / z
yxz, yzx p sxxz, xzx zxx, zzz
p p
zyy s ps s
29
2.3.3 Experimental consideration
The MSHG response depends not only on the geometry of
magnetization in relation to the incident light and sample plane, but also on
the polarization state of incident and outgoing laser beam. Table 2.2
categorizes the non-zero susceptibility tensor components obtained for
different polarization configurations. For the MSHG measurements on
PZT/LSMO heterostructure, the s-in s-out polarization configuration is used.
The odd susceptibility tensor element is xyyy(Mx) as a function of
magnetization along longitudinal direction, the even term comes from xzyy as
we rotate the detection analyzer 5 degree away from perfect s-out to
introduce a non-magnetic background for improving the magnetic contrast
(see below). The second harmonic signal comes from electrical polarization of
PZT mainly contributing to the term Xzzz and is eliminated in this optical
polarization configuration.
The measured intensity in a fixed experimental geometry with opposite
magnetization direction can, in general, be written as a sum of effective
tensor components:
IshH 2 co) oc \XeffW ± Xeff(~)\2 (2.15)
where T e //(+) and * e / / (_) are the linear combinations of even and odd tensor
components with Fresnel coefficients a iJk\
X e f f ^ l i , j . k & i j k X i j k (2-16)
30
The magnetic contrast A for MSHG measurement is defined as
where l(±M) is the signal intensity received by the detector when the
macroscopic magnetization M is aligned with external magnetic field in either
positive or negative magnetic field direction (Fig. 2.9). This contrast is
normalized to the total SHG intensity, and does not depend on the intensity of
the fundamental light. Moreover, the magnetic contrast A obtained by MSHG
is linear with M in the probed volume, if magnetic component is much smaller
than non-magnetic component in signal intensity as in our case [28].
0.36(+M)
- - I (-M)
0.28
-1000 0 1000Magnetic Field (Oe)
Figure 2.9 MSHG hysteresis loop from PZT/LSM O thin films indicating l(-M) and
l(+M) used to determine the magnetic contrast A from Eq. (2.17).
31
2.3.4 Experimental set-up
MSHG experiments are performed with a Ti:sapphire laser (RegA-
9000 amplifier system) generating 200 femtosecond pulses with 4 pJ energy
at 250 kHz repetition rate and 800 nm wavelength. The attenuated laser
beam (50 mW) with s-polarization is focused to a 200-pm diameter spot on
the sample at an angle of incidence of 40°. A small S-polarized MSHG signal
(400nm) is generated in the direction of the reflected laser beam, and is
detected with a high signal-to-noise ratio photomultiplier produced by
Hamamatsu company. Filtering with prism is required to separate the MSHG
light from the fundamental laser beam. The detection scheme A and B are for
MSHG and MOKE measurements respectively.
Red-psfiltei Prism
Lens Lens Photodiode
AnalyzesPolarizer
Figure 2.10 Schematic set-up of MSHG (A) and MOKE (B) measurement.
32
Chapter 3
Time resolved studies of spin dynamics in FePt thin flims
The coherent spin precession in L1o ordered FePt thin films is
investigated with TRMOKE technique. The perpendicular magnetic anisotropy
Ku and intrinsic Gilbert damping oq of the material systems are derived from
spin dynamics studies. Furthermore, the quadratic dependence of Ku and ao
on spin-orbit coupling strength £ is rigorously observed for the first time,
where £ is continuously controlled through chemical substitution of Pt with Pd
element. These results are consistent with first-principle calculations, which
also show that other leading physical parameters remain unchanged through
doping.
Moreover, a linear correlation between do and electron scattering rate
1/re is experimentally observed through modulating the anti-site disorder c in
the L10 ordered FePt thin films. The variation of c mainly affects re, while
other leading parameters remain unchanged. Moreover, as c decreases, the
perpendicular magnetic anisotropy increases and the Landau g factor exhibits
a negative shift due to an increase in orbital momentum anisotropy.
33
3.1 Introduction
Ultrafast magnetization precessional switching in ferromagnets utilizing
magnetic field pulses, spin polarized currents, and ultrafast laser pulses [29-
33] is currently a popular topic due to its importance in magnetic information
storage and spintronic applications. The uniform magnetization precession
can be well modeled with LLG equation, where the Gilbert damping
parameter a determines the spin relaxation time [34] and is crucial for device
performance [9-12]. The extrinsic Gilbert damping is due to nonlocal spin
relaxation, such as spin pumping and two-magnon scattering, which can be
tuned by artificial substrates, specially designed buffer and coverage layers
[35-41],
The intrinsic Gilbert damping a0 describes the energy flow rate from
spin to electronic orbital and phonon degrees of freedom through electron
scattering, and has been studied in various theoretical models [42-50]. The
breathing Fermi surface model [42] and torque-correlation model [43] based
on first-principle band-structure calculations qualitatively match or0 in soft
magnetic alloys, such as Fe,Co and Ni [51-56]. Moreover, contributions to do
can be categorized based on intraband and interband transitions [43, 46]. The
damping rate from intraband transitions scales linearly with the electron
relaxation time re and exhibits conductivity-like behavior. In contrast, the
damping rate from the interband transition is proportional to the electron
scattering rate 1/re, and consequently exhibits resistivity-like behavior.
Moreover, contributions of intraband (interband) transitions to damping are
34
predicted to play a dominant role in low (high) temperature regions, and be
proportional to £3 ( f 2).
Perpendicular magnetic anisotropy (PMA) defines the low-energy
orientation of the magnetization M (easy axis) normal to the film plane, as well
as the stability of M with respect to external fields, electric currents, and
temperature-induced fluctuations. Materials with large PMA are good
candidates for high-density memory and spintronics devices [8-10], as they
exhibit promising thermal stability and allow going beyond the
superparamagnetic effect, giving access to magnetic media with smaller
magnetic domains [7,8]. Moreover, PMA-based magnetic tunnel junctions
provide high tunneling magneto-resistance ratio [9-11], and low critical current
for spin-transfer-torque switching [12]. Considering these advantages,
tailoring the PMA of magnetic materials as well as elucidating its physical
origin becomes important for further technological advancements.
PMA can be related to the spin-orbit-coupling (SOC)-induced splitting
and shifting of electronic states that depend on the magnetization direction,
which results from the simultaneous occurrence of the spin polarization and
SOC [57-59]. An effective method to achieve significant PMA utilizes
appropriate combinations of 3d and heavier Ad, 5d elements, which merge
the large magnetic moment and magnetic stability of 3d TMs with the strong
SOC of Ad, 5d TMs [60-64], L10 ordered FePt material is one prominent
example [65-69]. In addition, Pt acquires a sizable spin polarization in contact
with Fe and contributes significantly to the PMA, due to its large SOC.
35
3.2 Samples and characterization
Figure 3.1 shows the schematic of atomic structure for the two series
of FePt thin films where the L10 ordered structure consists of alternate
stacking of Fe and Pt atomic planes along the face-center-tetragonal (fct)
[001] direction. In the first series (left, SOC dependent), continuous
substitution of Pt by Pd leads to a gradual decrease of £ in L1o ordered
FePd(i-X)Ptx alloy, while other leading physical parameters remain almost
unchanged. That is because the element Pt falls into the same group as Pd in
the periodic table, and £(Pt) is stronger than £(Pd). In the second series (right,
chemical order dependent), the anti-site occupation of atoms is controlled
with proper growth temperature, which mainly affects the electron scattering
rate 1/reand the hybridization effect between Fe and Pt atoms.
The two series of L10 FePt(Pd) alloy films are deposited on single
crystal MgO (001) substrates by magnetron sputtering. The FePt(Pd)
composite target is fabricated by placing small Pt (Pd) pieces on a Fe target.
The base pressure of the deposition system is 1.0*1 O'5 Pa and the Ar
pressure is 0.35 Pa. During deposition, the substrates are kept at 7"g=620 °C
for the SOC dependent study and at different temperatures Tg for the
chemical order dependent study, and the rate of deposition is about 0.1 nm/s.
After deposition, the samples are annealed in situ at the same temperature as
for the growth Tg for 2 hours. The film thickness is determined by x-ray
reflectivity to be 17 ± 1 nm (22 ± 1 nm).
36
t • • • • • • •• • • • • • •• • • • • • » Fe
[001] • • • • • • • ^• • • • • • •Perfect L1(0) ordered Fe/Pt alloy
/
• • • • • • •• • • • • • § Fe• • • • • • • • • • • • •
• • • • • • •• • • • • • « Fe• • • • • • •• • • • • • • Pt
• • • • • • • • Pd • • • • • • •FePt(x)Pd(l-x)
modulating spin-orbital coupling Varied chemical order
Figure 3.1 Schematics of atomic structures for the two series of FePt thin films.
Structure characterization of FePt(Pd) samples is performed with X-ray
diffraction measurement with Cu Ka radiation, as presented in Fig. 3.2. Only
fct (001) and (002) peaks of FePt(Pd) are observed in the spectrum along
with other peaks from MgO substrate, which indicates the L1o ordering in the
FePt(Pd) alloys. The peak positions do not shift with different Tg or Pd doping
level (1-x), which indicates that the lattice constant varies by less than 1.0
percent and tetragonal distortion of lattice is not affected by either doping or
modulation of Tg. The anti-site disorder percentage is derived from
°Q2' meas\ /omeas (3.1)('001)' 002' mi
37
where S is the degree of chemical order, /001 and /002 are integrated
intensities of fct (001) and (002) peaks, and (1 - c ) is the probability of the
correct site occupation for either Fe or Pt atoms in such L l0 ordered alloy
system [70]. U o o i / I oo2 )c a ic *s calculated to be 2.0 for the perfect ordered film
with thickness ranging from 11 to 49 nm [70], Thus for the Pd doped FePt thin
films, S is about 0.8, indicating a good suppression of disordering defects in
these films and about 90% of atoms stay at the correct sites. Moreover, c as a
function of Tgfor both 17-nm-thick and 22-nm-thick films is shown in Fig. 3.3.
Higher Tg leads to monotonous decrease of the anti-site disorder c in FePt
alloys, as depicted by the insets in Fig. 3.3.
38
X—1
x=0.5x=0.25
&toc<D-4—*
Qa : x fct (002)fct (001)
6040 5020 3020 (deg)
S
(/)c0coa:X
(b) fct (001)
Tg=720 °C f jfct (002)
u| Tg=680 °C P A
Tg=620 °C I I Aa Tg=580 °C
MgO J L1 1 i ......
20 4020
60
Figure 3.2 Structure characterization results of FePtxPd{1.x) thin films (a) and FePt
thin films with different Tg (b) by X-ray Diffraction measurements.
39
o
1 6
12
8
• • • • • • •• • • • • • * FeJ
• • • • • • •• • • • • • • Pti• • • • • • •
2 2 nm F e P t 17 nm F e P t
• • • • • • • •! • • • • • • f t •
- ' • • • • • • •Pt J• • • • • • • •
• • • • • • •
480 540 600 660Tg(°C)
720
Figure 3.3 Anti-site disorder percentage c in L10 ordered FePt thin films as a
function of growth temperature Tg.
40
"(a) X=1
“ (b)x=0.5
"(C)
— 0 . =ox=0- - 0 = 90
H (T)0 0006
0.0004
0.0002
0.0000 r
- 0.0002
-0.0004
-0 0006- 3 - 2 - 1 0 1 2 3
Magnetic field H (T)
Figure 3.4 Static magnetic hysteresis loops measured by vibrating sample
magnetometer (VSM) (a)-(c) for FePtxPd(i.X) thin films. The black (red) curves
represent out-of-plane (in-plane) results. VSM results for FePt thin films with
different Tg (d).
41
Figures 3.4(a)-(c) display the out-of-plane and in-plane magnetization
hysteresis loops of FePtxPd(i-X) thin films for x = 1.0, 0.5, and 0, respectively.
As shown in Fig. 3.4(a), for x =1 (L l0 FePt) the out-of-plane hysteresis loop is
almost square-shaped with coercivity Hc = 0.44 T; but it is hard to reach the
saturated magnetization with in-plane magnetic field, indicating the
establishment of high PMA. With decreasing x, Hc decreases due to the
reduced PMA. For x = 0 in Fig. 3.4(c), Hc approaches zero and the out-of
plane and in-plane hysteresis loops almost overlap with each other,
indicating weak PMA. The PMA therefore increases with increasing
concentration of Pt in FePd(i.X)Ptx alloy thin films. From the experiments, the
saturation magnetization Ms for all samples is 1100 emu/cm3 within 10%
relative error, close to the bulk value of L1o FePt [71].
Figure 3.4(d) displays the out-of-plane magnetization hysteresis loops
for 22-nm-thick films with Tg = 580 °C, 620 °C, and 680 °C, and the in-plane
hysteresis loop for the sample with Tg = 620 °C. The out-of-plane hysteresis
loops are almost square-shaped with coercivity Hc = 0.3 T; but it is difficult to
reach the saturated magnetization with in-plane magnetic field, indicating the
establishment of high perpendicular magnetic anisotropy Ku. From the
experiments, the saturation magnetization Ms for all samples is determined to
be 1100 emu/cm3 and remains unchanged as a function of growth
temperature and disorder c. Moreover, the magnetization is not fully saturated
with H = 2T applied along the easy axis, indicating the existence of multiple
magnetic domains at lower magnetic fields.
42
3.3 Magnetic anisotropy and damping in FePtxPd (1.x) thin films
Figure 3.5 shows TRMOKE results of FePd(i.X)Ptx films with various x
at H = 5 T. A uniform magnetization precession is excited as manifested by
the oscillatory Kerr signal 0k, while the magnetic damping is indicated by the
decay of precession amplitude as the time delay increases. As shown in Fig.
3.5, the measured Kerr signal can be well fitted by the following equation
0k = a+b*exp{-tAo)+A*exp(-tlT)s\r\(2jxft+q>) (3.2)
where parameters A, r, f and <p are the amplitude, magnetic relaxation time,
frequency, and initial phase of the magnetization precession, respectively.
Here, a, b, and to are related to the background signal owing to the slow
recovery process after fast demagnetization by laser pulse heating. It is well-
demonstrated in Fig. 3.5 that the spin precession frequency and magnetic
damping effect become larger and stronger for higher doping level x with the
same H.
In order to derive the PMA and ao for FePdi.xPtx samples at different
doping levels, magnetic field-dependent TRMOKE measurements are
performed. As shown by the TRMOKE results of x = 0.82 and x = 0.5 samples
in Fig. 3.6(a) and 3.6(b), spin precession frequency and relaxation time
increases as H increases. The field dependence of frequency can be
understood by the enhancement of the effective field due to larger H. As for
the damping, the strong magnetic field suppresses the extrinsic part of Gilbert
damping, such as caused by dephasing, and a is governed by intrinsic
43
contribution at higher magnetic fields. The H dependences of the precession
frequency f and relaxation time r for some typical samples are displayed in
Figs. 3.7(a)-3.7(b). W e note that f can be widely tuned from 15 GHz to 340
GHz by varying the magnetic field and doping level, and r decreases by two
orders of magnitude when Pd atoms are all replaced by Pt ones and reaches
about 7.0 ps for x=1 (L1o FePt).
From the LLG equation with cr«1.0, the dispersion equation is
2TTf = ><H,H2)1/2, (3.3)
where H i= H c o s ( 0 h - 0 ) + / - /kc o s 20 and /-/2= /- /c o s (0 w-0 )+ H k C O s 2 0 , Hk=2KJMs-
4ttMs with uniaxial magnetic anisotropy constant Ku, gyromagnetic ratio y
and 9h = 45°. The equilibrium angular position G of the magnetization satisfies
the following equation sin20 = (2/-///-/K)sin(0H-0). The measured field
dependence of f can be well fitted by Eq. (3.3), as shown in Fig. 3.7(a). With
the measured Ms of 1100 emu/cm3, and the g factor fitted to be 2.1 ± 0.05 for
samples with different x, Ku is derived from the dispersion and displayed in
Fig. 3.8(a) as a function of doping level x. For a « 1 .0 , aro can be in principle
extracted by fitting the measured field dependence of t with t =2/aroY/(Hi + H2)
and using the fitted values of g and HK, as shown in Fig. 3.7(b). The deviation
between fitting and data points at lower magnetic field is due to the
appearance of extrinsic Gilbert damping, which mainly comes from the
dephasing dynamics caused by inhomogeneity of PMA in thin films, as the
magnon scattering is less effective for perpendicularly magnetized samples
44
[72], The measured «o exhibits a sensitive dependence on the chemical
substitution in Fig. 3.8(b).
X=1
x=0.82
x=0.5zj
x=0.25CD
=0.15
30 4010 20Delay t (ps)
Figure 3.5 TRM OKE data with different doping level of FePtxPd(l-x) under magnetic
field H= 5T . The solid lines are fitted curves.
45
x= 0.82 H= 5 T
H= 4 T
H= 3 TCD
H= 2 T
30 4010 20Delay t (ps)
zsCD
o'
20 40 60 80Delay t (ps)
Figure 3.6 TRM OKE results of FePt0.82Pd0.18 (a) and FePt0.5Pd0.5 (b) under different magnetic fields H. The solid lines are fitted curves.
* X= 0 5 H= 5 Tj / V \ A a a a a a a / w w
i/ V V v v ^ ^
H= 2 T
46
0.3
0.2N
XI -
X=1x=0.9x=0.75x=0.5x=00.0
0 2 4 6H (T)
0.6• x=0.9 ▲ x=0.75 ▼ x=0.5♦ x=0
0.3
0.04
0.000 2 4 6
H(T)
Figure 3.7 The dependence of precession frequency f (a) and lifetime x (b) on
external magnetic field H fo r typical samples. The solid lines refer to fitted results.
47
^ 4COE ,0^5)0
r -O 2T—zs
00.0 0.5 1.0
X
0.10
0.05
0.00
Figure 3.8 PMA (a) and a0 (b) as a function of doping level x of FePtxPd(1.x).
Mb)
1.00.50.0
n
- /— i
•■ ■ ■
I i . i
48
0.6
>d) 0.4
0.2
15>
10
5 5
CD3.0ZL
U) 1.5F
0.01.5
<0 1.0O
0.5
-(b)
(a)
- (c)
'(d)
0.0
■I
__L _
0.5x
FePt/Pd
■ ■
1.0
Figure 3.9 Spin-orbit coupling strength f (a), bandwidth tV (b), spin moment mson Fe
or Pt/Pd site (c) and measured c/a ratio (d) as a function of doping level x of
FePtxPd(i.x).
49
To investigate the spin-orbit interaction in FePdi-xPtx doping system, ab
initio density functional calculations are performed to quantitatively determine
the key physical parameters [73, 74], The calculated f , bandwidth W and
magnetic moment ms at Fe or Pd/Pt site as a function of doping level x is
derived and depicted in Figs. 3.9(a)-3.9(c). Figure 3.9(a) shows that f at
Pd/Pt site changes from 0.19 to 0.57 eV when x varies from 0 to 1.0. For Pt,
Pd, and Fe atoms, £ is 0.6, 0.20, and 0.06 eV [73, 74], respectively, and
therefore the effect of Fe atoms is negligible compared with those of Pd and
Pt atoms. The change of spin moment and bandwidth are negligible when x is
tuned from 0 to 1, as shown in Fig. 3.9(b) and 3.9(c). The Pt/Pd atom
acquires spin polarization around 0.27 (pe/atom) due to the proximity effect.
Moreover, the lattice constant remains unchanged for different x from both
XRD results and theoretical structure analysis of L l0 FePdPt alloys, and c/a
ratio is around 0.94 as shown in Fig. 3(d). Furthermore, the density of states
at Fermi level D(EF) is found to change from 2.55 to 2.39 per atom per eV for
x varying from 0 to 1.0. Standard electron-phonon scattering calculations are
performed with the phonon Debye model and static limit of Lindhard’s
dielectric function. The electron-phonon scattering rate 1/re varies from 1.34
to 1.33 ps"1 for x changing from 0 to 1.0. Since other leading parameters
remain unchanged, the variation of PMA and ar0are mainly controlled by SOC
strength through different doping level x.
Figure 3.10 (a) displays the measured values of ar0as a function of£,
which agrees with the calculated values from first principles. The results
50
rigorously prove the theoretical prediction of the £2 scaling of ato for the first
time, indicating that ar0 at high temperatures is mainly caused by interband
contribution. Moreover, the electronic-scattering-based model of FM
relaxation is therefore proven to be applicable for ao in L10 FePdPt ternary
alloys.
Figure 3.10(b) shows the £ dependence of PMA derived from
TRMOKE measurements (black squares). The magnetic anisotropy arises
from second-order energy correction of SOC in the perturbation treatment.
Therefore, the PMA is roughly proportional to both f and orbital magnetic
moment when f is smaller than the exchange splitting. Since the orbital
moment is also proportional to ( within perturbation theory; the enhanced
PMA at higher x is attributed to a larger £ of Pt atoms compared with that of
Pd atoms and a quadratic scaling law of PMA with f is expected, as observed
from the experimental results. Moreover, first-principle calculations are
performed [74] and the calculated values of PMA are larger than the
measured ones at higher x (red spots in Fig. 3.10(b)). Several reasons exist
for the discrepancy, such as slight differences in the lattice structure between
experiment and theory, and thermal fluctuations and chemical disorder are
not accounted for in the theoretical calculations. Hence, there is potential for
further enhancing PMA by SOC tuning.
One approach to enhance the SOC dependence of PMA is increasing
the chemical order of FePdi.xPtx alloys, as the large PMA values in the
theoretical study are based on an ideally ordered lattice. To estimate the
51
influence of disorder, we perform TRMOKE experiments on the samples with
higher chemical ordering S = 0.85, where the correct site occupation of atoms
is increased to 92% from rFe = r Pt Pd) = 90% (S = 0.8). The measured PMA
values, indicated by the blue rhombus symbols in Fig. 3.10(b), show a slight
increase for x=0.5 (£ = 0.26) and a noticeable enhancement for x=1 (£ = 0.57).
This indicates that the chemical ordering of FePdi-xPtx alloys has a strong
effect on PMA by spin-orbit interaction, as the disordering can affect the
electronic band structure [75]. Also at lower temperature (20K) the PMA
values of FePdi.xPtx alloys can be more enhanced by SOC strength due to the
suppression of thermal spin fluctuation, as shown in Table 3.1. Furthermore,
other approaches such as modulation of tetragonal distortion and chemical
substitution of 3d transition metal elements in alloys may also be utilized for
SOC tuning of PMA.
52
E.o13>Q)
O
13
10■ measured • calculated8
6♦ chem ordered
4
2
0
0.0 0.3 0.6SOC (ev)
measured0.10
s° 0.05
0.00
0.0 0.3 0.6SOC $ (ev)
Figure 3.10 Measured and calculated cro (a) and PMA (b) as a function of SOC
strength. The solid line is a quadratic fit obtained from first principle calculations.
53
FePto.5Pdo.5
Temperature (K) 9 Hk(T) Ku (107erg/cm3)
20 2.13 2.17 2.07
200 2.10 2.08 1.914
FePto.7 5 Pdo.25
Temperature (K) 9 « k (T) Ku (107erg/cm3)
20 2.20 3.96 3.097
200 2.18 3.80 2.849
FePt
Temperature (K) 9 Hk (T) Ku (107erg/cm3)
20 1.97 7.30 5.012
200 2.00 6.96 4.587
Table 3.1 Landau g factors and Ku derived from TRMOKE measurements for
different doping level x at 20K and 200K.
54
3.4 Magnetic anisotropy and damping in FePt thin films with
different chemical order
Figure 3.11(a) shows the TRMOKE results of FePt thin films with
various c at H = 6.5 T. The uniform magnetization precession is demonstrated
by the oscillatory Kerr signals 0K, while the magnetic damping is indicated by
the decaying precession amplitude as the time delay increases. The
measured 6k can be well fitted by Eq. (3.2). It is well demonstrated in Fig.
3.11(a) that the spin precession frequency is larger and magnetic damping
effect becomes weaker for lower c with the same H.
In order to obtain the PMA and or for FePt samples with different c,
magnetic field (H) dependent TRMOKE measurements are performed. Figure
3.11(b) shows the measured results of the 22-nm-thick sample with Tg = 680
°C. It can be seen clearly that the precession period and relaxation time vary
as H increases. The fitted fas a function of H for different c are plotted in Fig.
3.12(a). We note that f can be tuned from 240 GHz to 335 GHz by varying H
and x. The measured H dependence of f can be well fitted by Eq. (3.3), as
shown in Fig. 3.12(a), and we thus obtain Ku and g. Ku decreases from 6.0 to
3.2 (107 erg/cm3) and the Landau g factor also displays a shift from 1.87 to
2.24 as the anti-site disorder percentage c increases, as shown in Figs.
3.13(a)-3.13(b).
55
13TO,*<D
(a)c= 4.0%
10
c= 8.6%
c= 10.5%
c= 13.8%
20 30t(ps)
c= 4.0% H= 6 T
H= 5 T=3d
H = 4 T
H= 3 T
20 30t(ps)
Figure 3.11 TRMOKE data from 22-nm thick FePt films, grown at different Tg with
applied magnetic field H -5T (a), and with Tg =720 °C measured at different H (b).
The solid lines refer to fitted results.
56
0.35c= 4.0% c= 8.6% c— 10.5% c= 13.8%0.30
NII -- 0.25
0.202 64
H (T)
▼ c= 13.8% a c= 10.5% • c= 8.6%■ c= 4.0%
0.20
0.15
4)
0.10
0.05
2 4 6H (T)
Figure 3.12 The dependence of spin precession frequency f (a) and effective Gilbert
damping aeff (b) on H for FePt thin films with different c. The solid lines are fitted
results.
57
EoE>a>
h-o
=3
(a)
■ 22nm FePt - 17nm FePt
0 6 12 Anti-site percentage c
18
2.2
oo£ 2.0o>
6 12 180Anti-site percentage c
Figure 3.13 PMA Ku (a) and Landau g factor (b) as a function of anti-site defect
concentration c for FePt films with thicknesses of 17 nm and 22 nm.
58
Figure 3.13(a) shows that the Ku maintains high values from 3.2 to 6.0
(107 erg/cm3), and gradually increases with smaller c, which is consistent with
other experiments, as well as the theories [76-79]. The Ku in FePt alloys
results from the simultaneous occurrence of the spin polarization and large
SOC [74], The smaller c indicates that more Pt atoms become nearest
neighbors of Fe, which results in stronger hybridization between Fe and Pt
atoms. Consequently, Pt acquires larger spin polarization and orbital moment
due to the Fe-Pt hybridization in the sample with higher chemical order, and
contribute significantly to the Ku. As a result, the orbital momentum anisotropy
[78, 79] and hence Ku increase with decreasing c.
Figure 3.13(b) reveals a gradual decrease of Landau g factor with
smaller c. The g factor describes the proportionality of angular momentum
and magnetic moment for the individual spins, which also affects the dynamic
response of a magnetic film. In itinerant electron systems, the g factor may be
written as
(3.4)y e <S'>+</-> v '
where ps (ul) denotes the magnetic moment from the spin (orbit)
component. For a symmetric crystal lattice, the orbital motion of d electron is
quenched by crystal field effect, i.e., <L> = 0. Nevertheless, the orbital
contribution to the magnetic moment is nonzero, thus the g factor is equal to
2*(1+ piJps) and is typically greater than two in itinerant electron system.
However, as c decreases, the enhanced hybridization between Fe and Pt
59
restores the orbital momentum due to the large SOC strength of Pt [78], and
raises the orbital momentum anisotropy [78, 79]. This would lead to <L> * 0,
and s f 2 * ( 1 - g J fJ s ) , indicating a negative shift of g factor relative to the value
of two. Such negative shift of g factor is also observed at surface or interface,
where the orbital momentum is not entirely quenched due to the broken
symmetry [80].
Using the fitted values of r, we determine a with
a , „ = a„ + a „ = (3.4)
where arex is the extrinsic contribution to Gilbert damping. As shown in Fig.
3.12(b), the value of aeff gradually decreases with higher H and saturates at
higher fields. The decreasing trend of creff with H is attributed to the
suppression of dephasing dynamics among magnetic domains, since multiple
domains exist at low fields as indicated in Fig. 3.12 (b) and the magnon-
magnon scattering is less effective for perpendicularly magnetized samples
[72]. The saturation values of areff at higher fields are therefore used to
approximate the intrinsic Gilbert damping ao [73]. Figure 3.14(a) shows cto as
a function of c at 7=200K.
60
0.21
o>■£= 0.14Q .E(0■O■Ca>
S 007
(a)22nm FePt 17nm FePt
0 6 12 18 Anti-site percentage c
0.09
0.06o>
0.03
22nm FePt 17nm FePt
0 6 12 18 Anti-site percentage c
Figure 3.14 a0 (a), and or</g (b) as a function of anti-site defect concentration c for
FePt films with thicknesses of 17 nm and 22 nm. The solid line refers to least square
linear fit.
61
The key finding here is that or0 gradually increases with larger c. Since
the scattering rate 1/re is enhanced with more impurity scattering sites c
according to scattering theory [48-50], the positive correlation between aro and
1/je qualitatively matches the resistivity-like behavior where or0 is governed by
interband transitions [43, 46]. In the spin-orbit coupling torque correlation
model for interband transitions [43, 46],
a„ « g ( f /M 0 2A« (3.5)
where W is the d-band width, and f the spin-orbit coupling strength. It is
demonstrated that anti-site disorder in L l0 ordered alloys smoothens the
density of states [81]. The calculated D(Ep) increases less than 5% as c
increases, and W remains almost the same [75], when 0<c<15% as in our
case. In our previous work [82], we investigated the anomalous Hall
conductivity of L1o ordered FePt films with different ordering. We found that
the resistivity independent term (bo) of anomalous Hall conductivity remains
almost unchanged with x when x is small, indicating that the variance of spin-
orbit coupling strength £ in our samples is negligible. The g factor increases
by 1 9 % with larger c as shown in Fig. 3(a), due to the modulation of orbital
momentum anisotropy. The lattice distortion and Ms remain almost
unchanged as demonstrated by structure characterization and VSM
measurements. A similar trend of aro with c is observed for samples with two
different thicknesses, which indicates that surface and strain effects on
damping can be ruled out. Therefore, we attribute the significant increase of
62
aQ by more than a factor of three, as revealed in Fig. 3.14(a), to the
enhancement of 1/re as a result of the increasing c. To separate the effect of
g on the damping, we further plot aotg as a function of c in Fig. 3.14(b). We
observe approximately a linear correlation between the two variables. The
exchange of different elements in L10 ordered alloy film leads to scattering of
itinerant electrons through local spin dependent exchange potentials. Since
the cross-section and strength of individual scattering events remain
unchanged in weak scattering regime [48], the linear correlation indicates that
do is proportional to 1/re, where the damping process is dominated by
interband contribution. The damping process can be considered roughly as
the decay of a uniform mode magnon into an electron spin-flip excitation.
Therefore, the anti-site disorder works as spin-flip scattering center for
itinerant electrons transferring spin angular momentum to the lattice via SOC.
Moreover, complete suppression of the anti-site defects might lead to a
remnant cro, where electron is mainly scattered by phonon instead of impurity.
In order to check for the intraband contribution to cro, which may become
significant with low 1/re, TRMOKE measurements are carried out for the 17-
nm-thick FePt film with the lowest c=3% at low temperature. Figure 3.15(a)
shows that the frequency and decay rate of coherent spin precession at H=5T
varies slightly with temperature (from 20 K to 200 K). The field dependent
frequency and decay rate at 20 K, as shown in Fig. 3.15(b), are analyzed to
obtain a0. It turns out that a0 remains almost unchanged (from 0.053 ± 0.013
to 0.054 ± 0.013) when temperature decreases from 200 K to 20 K, despite a
63
significant change in the electron-phonon scattering rate [43, 46]. The
temperature independent behavior of cro is due to the fact that electron-
impurity scattering dominates the scattering events. In Chapter 3.3, we
calculated the electron-phonon scattering rate from first principles to be
approximately 1.33 ps'1 for FePt at 200 K. The electron-impurity scattering
rate must be considerably higher, consistent with the values of electron
scattering rate from Ref. [46] where the interband contribution to cro dominates.
Moreover, the weak temperature dependence of cro indicates that 3% anti-site
disorder is still too high to observe the conductivity-like behavior for intraband
contribution to cr0, which is consistent with theoretical prediction [49]. Further
investigations at low temperature with fewer impurities are necessary to
elucidate the relationship between cr0 and 1/re governed by intraband
transitions.
64
0.3 (a)
N 0.2
0.1 - •
0
Frequency Decay rate
100 T (K)
0.3
0.2 £
0.1
200
0.30.3
■ Frequency • Decay rate0.2
NXI—
3 64 5H (T)
Figure 3.15 The dependence of frequency f and decay rate 1/ron temperature T
with applied field H =5T (a), and as a function of H at 20 K (b). The solid lines refer to
fitted results.
65
Chapter 4
Time resolved spin precession in Fe/CoO heterostructure
Optical excitation of spin precession is investigated in Fe/CoO (FM-
AFM) exchange coupled heterostructure with TRMOKE technique. Photo
excited charge transfer processes in AFM CoO layer create a strong transient
exchange torque on FM Fe layer through FM-AFM exchange coupling, where
the efficiency of spin precession excitation is significantly higher and the
recovery is notably faster than the demagnetization procedure. The
precession amplitude peaks around room temperature and with external
magnetic field competitive to the magnetic anisotropy field, indicating that this
efficient excitation mechanism originates from the modulation of the uniaxial
magnetic anisotropy Ku induced by the FM-AFM exchange coupling. Our
results will help promote the development of low energy consumption
magnetic device concepts for fast spin manipulation at room temperature.
66
4.1 Introduction
Control of coherent spin precession in ferromagnets is currently a
popular topic due to its importance in magnetic recording and spintronic
devices [83-92]. Recently, optical excitation of coherent spin precession
utilizing ultrafast laser pulses provides novel approaches for fast spin
manipulation [83-89] and generation of ultrafast spin currents [86]. The
precessions are usually triggered by the transient torque on magnetization
generated via spin transfer effect [13, 90], demagnetization and modulation of
magnetic anisotropy [93-95]. The rapid demagnetization may be caused by
dissipation of spin angular momentum through electron-lattice interaction [83,
84], or transport of super-diffusive spin-polarized current [85], whereas the
anisotropy constant modulation may be caused by the alteration of spin-orbit
interactions due to hot electrons and electron-phonon thermalization [93],
creation of itinerant carriers [94], or coherent acoustic phonons [95].
Despite the different excitation mechanisms, most scenarios utilize the
optical absorption followed by a rapid raise in temperature and hence require
intense laser pulses. The long cooling time and excessive heating
considerably limit the push for fast and energy-efficient spin manipulation. For
instance, the demagnetization procedure usually requires intense optical
pluses to generate the non-equilibrium hot electrons [84] and thermodynamic
gradients [86]. In exchange-biased magnetic structures, the instant and
intense laser heating modifies the exchange bias, destroys pinning of the FM
spins at the interface, and thereby triggers the rotation of the magnetization to
67
the new direction [96]. But still, the thermal reduction of exchange bias is
followed by a few-hundred-picosecond recovery through heat diffusion.
The search for non-thermal excitation mechanisms motivates
extensive research to overcome the disadvantages of thermal ones [87, 88,
97-101]. The main idea is to utilize the interaction between magnetization and
photo-excited carriers that are selectively optical-pumped, where the
recombination time of photo-carriers is much shorter than the heat diffusion
process. For instance, optical spin-transfer [87] and spin-orbit [88] torques
are observed recently in FM semiconducting GaMnAs, which however usually
happens at low temperature. An alternative approach is through FM-AFM
exchange coupling, as small modulation of the exchange coupling strength
might lead to notable changes in magnetic properties [102, 103]. It is recently
demonstrated that short laser pulses can non-thermally introduce spin
reorientation and dynamics in AFM materials much faster than in FM
materials [104, 105]. But the question is still open whether it is possible to
drive FM at the speed of AFM through FM-AFM exchange across
heterostructure interface.
68
4.2 Samples and Characterizations
The Fe/CoO thin films are grown on the MgO(OOI) substrate in an
ultrahigh vacuum chamber. The MgO(OOI) substrate with a miscut angle less
than 0.5° is prepared by annealing at 600 °C for 30 min. The CoO thin films
are grown by a reactive deposition of Co under an oxygen pressure of
2 * 10-6 Torr. The thickness of CoO layer is 3 nm. Then, a 4 nm Fe film is
epitaxially grown on top of the CoO film at room temperature (RT). All of the
samples are covered by a 3-nm-thick MgO protection layer. The high quality
epitaxial growth of CoO film and Fe film can be proven by the reflection high
energy electron diffraction (RHEED) patterns [102], and the epitaxial relation
is CoO[110]//Fe[100],
Longitudinal MOKE measurements indicate a negligible exchange bias
(/-/e) in Fe/CoO from 80 K to above RT, as shown in Figs. 4.1 and 4.2. The
absence of exchange bias is likely due to a very small number of
uncompensated AFM spins at the smooth interface. However, the field
cooling leads to the alignment of AFM spins along CoO [110] directions,
which is collinear to Fe [100]. Due to the exchange coupling of the Fe
magnetization with the AFM spins, a uniaxial anisotropy appears below 7n
(discussed below), and the coercivity (Hc) increases with decreasing
temperature down to 130 K, below which Hc decreases because less AFM
spins are dragged to switch with FM spins [103]. As shown in Fig. 4.1, the
hysteresis loop is almost square for the field along the easy axis and is much
sharper at 80 K than 330 K; however, it is hard to reach saturation
69
magnetization at the largest field, restricted by our electromagnet, along the
hard axis perpendicular to the cooling field. Such results indicate a well-
defined easy axis and homogeneous anisotropy of the sample.
C Om mmm
c
■
-Qco
coco>
0 )
CD*
Figure 4.1 Longitudinal hysteresis loops for magnetic field along uniaxial easy axis
[100] and hard axis [010] at 82 K and 330 K.
8 2 K
H / / [ 1 0 0 ]
3 3 0 K
H / / [ 1 0 0 ]
- 1.0 - 0.5 0.0 0.5 1.0 Field(KOe)
70
600
^ 4 0 0
0
£ 2 0 0o 6
o
1 o
T -------------- 1-------------- 1-------------- ,---------------1-
HeHe
90 180 270 360
T(K)
Figure 4 .2 Temperature dependence of coercivity (Hc) and exchange biasing field
(He) obtained from easy axis loops.
71
4.3 Magnetic anisotropy
Figure 4,3 presents the TRMOKE results obtained from Fe/CoO thin
films with different magnetic field H along the Fe [110] direction (in-plane hard
axis) at RT. The oscillation is fitted to a damped sine function,
Aexp(-f/r )cos(2-nft), with precession amplitude A, time delay t , decay rate
1/r , and precession frequency f . It is obvious that spin precession frequency
f and relaxation time r increases as H increases. Figure 4.4 displays f as a
function of H. By solving the LLG equation, the frequency f can be well fitted
with the following equation:
2 = y J [ H cos( S - f ) + H a ] [ H cos(<?- j ) + H p] , (4.1)
where
Ha =Hdeff- 2KuS'm2flM s + K , ( 2 - sin22 <j>)IMs,
H? = 2K,cos40Ms + 2KuCOs2flMs, (4.2)
with effective demagnetization field Hdeff=4nMs+2K±/Ms, saturated
magnetization Ms, and gyromagnetic ratio y. </> and 5 are the angles of in-plane
equilibrium magnetization and H with respect to the Fe axis [100]. Ku, Ki and
Kx are the in-plane uniaxial, crystalline cubic, and out-of-plane magnetic
anisotropies, respectively. The <f> is derived from searching the minimum of
magnetic free energy. Therefore, the values of magnetic anisotropies at RT
are derived as: Hdeff= 2.0 kOe, KJMS=205 Oe and K1IMS=290 Oe.
72
•— 2112■ 1848■ 1320— 658 ► 461■ 330
264• 198— 132 >— 66
V
* *
05
— minus132
0.0 0.1 0.2- 0.1 0.3 0.4 0.5Delay time t (ns)
Figure 4.3 TRM OKE results of Fe/CoO thin films with different magnetic field H along
the Fe [110] direction (in-plane hard axis) at RT.
73
20
NXo
1000 2000 3000-1000 0
H (O e)
Figure 4 .4 spin precession frequency fa s a function of magnetic field H.
74
In order to study the uniaxial magnetic anisotropy Ku induced by FM-
AFM exchange coupling between Fe and CoO layer, temperature dependent
TRMOKE measurements are carried out. Figures 4.5 and 4.6 display the
TRMOKE results with magnetic field H= 3 kOe and 1 kOe applied along Fe
[11 yfl ] direction. The magnetization precession is faster at lower temperature,
due to the enhanced effective magnetic field Heff by the increase of Ku. Figure
4.7 summarizes the precession frequency f as a function of temperature.
Since the geometry of applied magnetic field is changed, the dispersion
relationship between f and H deduced from LLG equation requires
modification. By using the method mentioned in Chapter 2.2, we can deduce
the torque equations as follows:
\~dLt~ = = m<p [ ^ 1(2sin40sin2< cos2^ + 2sin20cos20 — sin20(sin4<p +
cos4<p) + 4sin20sin2<»cos2<p — cos20) + — (sin20sin2<p — cos2<p) +Ms
2cos26 + 2nMs) — ^ (sinBcoscp + sinQsin<p + a/2c o s 0)] ,
idmjp _ _ m&\^ (2 s m 2Bsin2(pcos2(pcos20 + 4sin29cos29sin2<pcos2(p +Y a t M s
cos40) + ^ Lsin2<pcos26 — 2cos29 + 2nMs j +^{sin9cos(p + sinBsincp +
yftcosB)] (4.3)
where 9 (<p) the angle between magnetization and Fe [001] (its projection on
(001) plane and Fe [100]) direction in the spherical coordinate system. By
treating the mg and with wave form as in Chapter 2.2, one can derive:
2kf = yJW Jh (4.4)
75
3nj,lccg»to
a)*ca)'</>cCO
j * ‘ . * •* *''** .*—*7 \ V a ^ - ^ wr \
»*S4
A U * U * ~ ^s .-. •. ; * ♦/« y ***
* * 4 « 4 4 ^ ^ _
-t._/\ yv y \ v \ - / ,%-
300K290K280K270K260K250K240K230K210K190K150K100K50K10K
0.0 0.1 0.2 0.3 0.4 0.5
t(n s )
Figure 4.5 Temperature dependent TRM OKE results with magnetic field H= 3 kOe.
76
Tran
sien
t Ke
rr si
gnal
(a.
u.)
• • ■. / ......
300K290K280K270K260K250K240K230K210K190K150K100K50K10K
Figure 4 .6 Temperature dependent TRM OKE results with magnetic field H= 1 kOe.
77
35
30
_ 25NXo
20
15
Figure 4.7 spin precession frequency f as a function of temperature T.
«
H = 1 kOe H = 3 kOe
0 50 100 150 200 250 300
T ( K)
78
The next step is to determine the equilibrium direction of magnetization,
i.e. the values of 0 and (p. The magnetic free energy F can be written as:
F = — (sin4 0sin2 2d> + sin2 20) + K u sin2 ^sin2 0 + K, cos2 0 + 2 /r(M cos#)24 (4.5)
- (sin 0 cos <j> + sin 0 sin <f> + 42 cos 0)
Along the equilibrium direction of magnetization, F reaches the
d F dFminimum, so that — = 0 and — = 0. Therefore, 0 and q> can be expressed
as a function of anisotropy fields and H. By fitting the field dependence of f,
one can obtain the equilibrium direction of magnetization as well as the
anisotropy fields.
f 25
O.
15 -
100 150Temperature (K)
200 250 300
Figure 4.8 q> as a function of temperature (green=3 kOe, blue=1 kOe).
79
thet
a (d
egre
e)
88.5
87.5
86.5
85.5
84.5
100 150Temperature (K)
200 250 300
Figure 4.9 G as a function of temperature (green=3 kOe, blue=1 kOe).
80
2.0
0O
Ku/Ms
0.5
0.00 100 200 300
T (K)
Figure 4.10 Uniaxial magnetic anisotropy field KJM S as a function of temperature T.
Figure 4.10 shows the dependence of uniaxial magnetic anisotropy
field KJMs on the temperature. Ku increases as T decreases. This indicates
that its origin is the AFM ordering of CoO layer. The FM-AFM exchange
coupling between Fe and CoO layer establishes an extra preference of
magnetization alignment in Fe, where the FM spins favor aligning parallel with
the frozen AFM spins. As the number of frozen AFM spins in CoO is
increased with lower T, the FM/-AFM exchange coupling is enhanced, which
leads to an increase of Ku.
81
4.4 Ultrafast spin exchange torque via photoexcited charge
transfer processes
TRMOKE measurements are performed in a pump-probe setup, where
the intensity ratio of the pump to probe pulses is set to be about 6:1. The
probe (A=800 nm) utilizes the MOKE technique with crossed polarizers and
incident angle around 40° to investigate the transient magnetic state change
along longitudinal and polar directions. Two strategies of pump are employed
in the measurements to investigate the optical excitation mechanisms: First,
the more intense pump pulses (A=800 nm, 3.1 mJ/cm2) are used to modulate
the FM order of Fe layer as shown in Fig. 4.11 (left), since the CoO layer is
almost transparent to 800-nm light [106]; Second, the weaker pump pulses
(A=400 nm, 0.16 mJ/cm2) are utilized to mainly affect the AFM order of CoO
as depicted in Fig. 4.11 (right). All TRMOKE measurements are performed
after field cooling the sample.
The TRMOKE result from Fe/MgO heterostructure is displayed in Fig.
4.12 (black squares) with pump pulse fluence 3.1 mJ/cm2 and magnetic field
H= 2 kOe at room temperature. The sudden rise and decay of Kerr signal
corresponds to the demagnetization procedure as described by the three-
temperature model [107, 108]. The rise in electron temperature by
demagnetization is first partially relaxed to the lattice within several
picoseconds. Afterwards, the heat of electron, spin and lattice diffuses to the
surrounding over >500 ps, which slowly recovers the system to its ground
82
state. Meanwhile, the magnetization starts to precess back to the previous
equilibrium direction in a damped circling way described by LLG equation. As
a result, the measured Kerr signal can be well fitted by the following equation
6|<= a*exp(-t/?o)+A*exp(-f/ r)sin(2ir ft+q>), (4.6)
where parameters A, r, f and <p are the amplitude, magnetic relaxation time,
frequency, and initial phase of the magnetization precession mainly along the
polar direction, respectively. Here, a and to are related to the background
signal owing to the slow recovery of magnetization after fast demagnetization
by the pump pulses, which happens mainly along the longitudinal direction.
The amplitude a qualitatively describes how far the FM system is driven away
from the ground state, and the amplitude ratio A/a can be used to qualitatively
estimate the efficiency of optically-driven coherent spin precession, which is
around 0.16 for Fe/MgO (black squares in Fig. 4.12).
Figure 4.12 (red circles) shows the TRMOKE result from Fe/CoO(3
nm)/MgO structure with the same excitation condition as the measurement on
Fe/MgO (black squares). We note that the background amplitude a remains
unchanged between the two measurements, as shown in Fig. 4.12, while in
contrast the amplitude of coherent spin precession A is enhanced with A/a =
0.85 for Fe/CoO/MgO heterostructure (red circles). This shows that the AFM
CoO layer improves the efficiency of optically excited coherent spin
precession. In Fe/MgO, A can be understood as Ai+Am, where A and Am
denote the contributions from modulation of magnetic anisotropy field Ha
83
through reduction of cubic magnetic anisotropy constant Ki and
magnetization Ms, respectively. Whereas in Fe/CoO, there exists another
contribution Au, which originates from modulation of the uniaxial magnetic
anisotropy constant Ku induced by FM-AFM exchange coupling between Fe
and CoO layers. Since Ku is more sensitive to temperature rise compared with
Ki [103], the modulation of Ku through laser pulse heating contributes
significantly to the enhancement of spin precession excitation.
The key finding here is that pronounced spin precession is still
observed with much lower pump pulse energy 0.16 mJ/cm2 at 400-nm
wavelength, as shown in Fig. 4.12 (blue triangles). Moreover, the TRMOKE
data reveal the absence of obvious demagnetization and slow recovery of
a Ms. Therefore, the rise of electron temperature Te in Fe layer through laser
pulse heating is negligible by the application of low pump pulse intensity.
Thus, the amplitude ratio A/a =8.35 is about two orders of magnitude larger
than from Fe/MgO with stronger pump pulse. Also, the instant pronounced
spin precession points towards an ultrafast non-thermal excitation process in
the AFM CoO layer as depicted in Fig. 4.11 (right) and discussed further
below. Furthermore, the magnetic relaxation time t decreases from 330 ps
(red circles) to 100 ps (blue triangles) in Fe/CoO with gentler pump pulses,
which is desirable for fast switching [109]. This is due to the fact that Gilbert
damping is enhanced at lower temperature in Fe/CoO exchange system [103].
84
800 nm 400 nm
®@®@@ ® @ ®
XX• • • •
FM Fe -
FM/AFM exchange
— AFM CoO —
*@@@@
i Co 3d
3 > 0 2 ph v ~ 3 . l e V
N(E)
Figure 4.11 Two different pump strategies to optically excite the coherent spin
precession, where the black arrows represent the spins of magnetic ions. The optical
charge transfer transition in CoO is depicted.
85
=3aj
roca'tok.fc
*
0 200 400Delay t(ps)
Figure 4.12 TRM OKE results from Fe/MgO (black squares) and Fe/CoO (red circles)
with pump pulse intensity 3.1 mJ/cm2, and Fe/CoO (blue triangles) with pump
intensity 0.16 mJ/cm2.
- Fe/MgO 3.1 mJ/cm2 • Fe/CoO 3.1 mJ/cm2 a Fe/CoO 0.16 mJ/cm2
• \ j l ,A=800nm
^ 'A A A A a aA=400 nm
86
The optical excitation of spin precession in Fe/CoO heterostructure can
be due to either the demagnetization procedure, or a modulation of magnetic
anisotropy constants, or spin transfer torque, or a combination thereof. From
the results described above, the fast demagnetization procedure is not
observed and the saturated magnetization Ms remains almost unchanged,
where AM scales down with lower a. Moreover, the modulation of magnetic
anisotropy constant Ki in Fe layer is reduced by the decrease of pump pulse
energy (i.e. Ai also contributes much less). Another possibility of the
enhanced Ala is the spin transfer induced torque. Even though the CoO
insulating layer eliminates the spin current transfer via spin pumping effect [9],
we have shown recently that the FM/AFM exchange establishes a spin
angular momentum transfer channel in Fe/CoO heterostructure, which slightly
increases with lower temperature [103]. To check this possibility, temperature
dependent TRMOKE measurements are carried out. Figures 4.5 and 4.6
present the TRMOKE results with pump pulse intensity 0.16 mJ/cm2 at
different temperatures T, where the absence of demagnetization peak is
observed ranging from low to room temperature. Figure 4.13 summarizes the
precession amplitude A as a function of temperature T with different magnetic
field H applied, and shows that A is peaked around 270K-290K, indicating
that spin angular momentum transfer is not the main contribution. Therefore,
the origin of this efficient excitation mechanism has to be the modulation of
magnetic anisotropy constant Ku as a result of variance in FM-AFM exchange
coupling.
87
H = 0.7 kOe H = 2 kOe20
< 10
01 0 0 200
T (K )300
Figure 4 .13 Measured and simulated spin precession amplitude A as a function of T.
88
T= 240K
T= 300K
T= 78K
00 1 2 3
H (kOe)
Figure 4.14 The spin precession amplitude A is plotted as a function of H at different
T, where solid lines represent simulated results.
89
The modulation of Ku also sheds light on the difference between the
two optical excitation strategies, 800-nm versus 400-nm wavelength pulses.
On one hand, the CoO layer is almost transparent to the light with A=800 nm,
as its bandgap is around 2.5 eV [106]. Therefore, the 800-nm pump pulses
mainly raise the electron temperature in FM Fe layer as depicted in Fig.
4.11 (left), and modulate the FM-AFM exchange coupling and hence Ku. This
leads to improvement of spin precession excitation efficiency as compared
with Fe/MgO thin film (Fig. 4.12 black squares and red circles). It is also
reported by Ju. et al. [96] that the heat in FM layer partially diffuses to the
adjunct AFM layer and slightly modulates the interface AFM order. On the
other hand, the FM order in Fe layer is much less modulated by the gentle
pump pulses with A=400 nm (3.1 eV), while there exist photon-excited
electron transitions in CoO layer as depicted in Fig. 4.11 (right). The near-gap
photons excite charge transfer transitions from O 2p band to Co 3d band that
is partially occupied with minority spins [110], which promotes the nearest-
neighbor FM exchange interaction and reduces the AFM ordering in CoO
[111], Since the coupling between Fe and CoO is a long-range interaction
[112], the variance of AFM order in CoO layer notably modulates the FM/AFM
exchange coupling and Ku, and thus leads to significant enhancement of the
excited spin precession (Fig. 4.12 blue triangles).
In order to derive the value of Ku, magnetic field dependent TRMOKE
measurements are carried out at different T. Figure 2(c) (black squares and
red circles) displays the spin precession frequency f for two different fields H
90
as a function of 7, where f increases as H increases. This can be understood
by the enhancement of the effective field due to larger H and is well fitted with
LLG equation to derive the magnetic anisotropies. It turns out that the cubic
magnetic anisotropy Ki originating from Fe layer remains almost unchanged,
while the uniaxial magnetic anisotropy Ku induced by the FM/AFM exchange
coupling between Fe and CoO layers increases significantly with decreasing
T, as shown by the blue triangles in Fig. 4.10. The derived KJMs behaves
similarly like f as a function of 7. The FM-AFM exchange coupling establishes
an extra preference of magnetization alignment in Fe, where the FM spins
favor aligning parallel with the frozen AFM spins. As the number of frozen
AFM spins in CoO is increased with lower T, the FM/AFM exchange coupling
is enhanced, which leads to an increase of Ku and hence increases f with
lower 7 as shown in Figs. 4.7 and 4.10.
To further elucidate the coherent spin precession excited by
modulation of anisotropy constant Ku, the relationship between precession
amplitude A and applied field H is investigated. Although the modulation of
anisotropy field Ha is independent of H, the transient torque on the
magnetization becomes stronger with increasing H as it pulls the
magnetization away from Ha, thus leading to larger precession angles at
higher fields. When the magnetization is nearly aligned along H at the field
strength stronger than Ha, the torque starts to saturate and the precession
angle decreases with increasing H. Thus, one can expect that the precession
amplitude reaches a maximum at the field approximately equal to the static
91
anisotropy field. Figure 4.14 shows the value of A as a function of H. We note
that at T = 300K the amplitude A (black squares) is first sharply enhanced
with increasing field and then saturates around H = 1 kOe and at last slowly
decreases, consistent with our expectation. Moreover, the simulation of A
versus H is carried out, where A is assumed to be proportional to the
equilibrium direction change by modulation of Ku (AKU) [113-115]. The AKU
and scaling factor s are parameterized at 300K. The black curve in Fig. 2(d)
corresponds to the simulation of A with a reduction of Ku about (165 Oe*Ms),
which agrees quite well with the measured relationship between A and H.
The simulation results also shed light on the evolution of precession
amplitude A with temperature T shown in Fig. 4.13. As T decreases, Ha is
enhanced by Ku, which requires stronger H to compete with Ha and the
generated transient torque is saturated. As a result, the peak in amplitude
peak is shifted to higher fields and becomes broader with lower T (Fig. 4.14).
The blue curve in Fig. 4.14 presents the simulated result at T=240K with s as
derived above and AKU parameterized equivalent to the effect of 20-K rise in 7.
The amplitude peak is around H - 1.8 kOe and the rise and decay of A due to
increasing H is less steep than at 7=300K. The shifting and broadening of the
amplitude peak is further enhanced at 78K as illustrated by the measured
results (red circles). Therefore, the evolution of amplitude A with temperature
7 in Fig. 4.13 can be interpreted as the shift of amplitude peak. Figure 4.13
shows the simulated amplitude A as a function of 7 for two different fields H
(dashed lines), which agrees quite well with the measured values. In this
92
simulation, AKU corresponds to a temperature rise of A7=20K, which is much
larger than AT «1K in CoO layer due to direct heating from optical absorption.
The photo-excited charge transfer processes essentially raise the
temperature of the CoO spin system, which effectively modulates the AFM
order of CoO.
In order to investigate the duration of this fast FM/AFM exchange
torque, simulations of real-time magnetization precession are carried out from
LLG equation. The observed pronounced coherent spin precession starts
right at t= 0, indicating a sudden change of equilibrium direction caused by
AKu. W e assume that the reduction of Ku (165 Oe*Ms) happens instantly at
t=0, followed by an exponential recovery process with time constant rr. Real
time simulation of magnetization precession is carried out with LLG equations
with time interval Af=0.2ps:
mz{t+At)IMs= mz(t)/Ms*exp(1-At/T) - my(t)l Ms *Y (Hcos(6-(p)+H^(t))*At
my(t+At)/Ms= my(t)/Ms*exp(1-Af/T) + mz(f+Af)*y (Hcos{6-<p)+hf(t))*At +(<p(Ku -
AKu(t+At))- <p(Ku -AKu(t))) (4.7)
where the first term represents the Gilbert damping effect and the second
term denotes the circling of spin precession. As for the my, there is an
additional term describing the change of equilibrium direction driven by
modulation of Ku.
As shown in Figs. 4.15 with rr = 40 ps, the simulated time evolution of
magnetization (red curve) agrees quite well with the probed Kerr signal (black
93
dots), since the TRMOKE signal is proportional to the polar component of
magnetization precession. Shorter or longer t> leads to mismatch in the
oscillation phase, as shown in Figs. 4.16-4.18. The fast recovery from photo
excited transitions might be due to the strong carrier-phonon interaction,
which results in non-radiative and phonon-assisted carrier recombination
[116]. Moreover, the recovery time of the instant photo-induced exchange
torque is much faster than the cooling time from demagnetization, which
might promote novel device concepts for fast spin manipulation.
step size = 0.8 ps xr = 40 ps £0.02
0.01
s 0.00o>
- 0.01AAT,
- 0.02
800 40Delay t (ps)
Figure 4.15 Measured Kerr signal change with 0.8 ps step size (black squares),
simulated polar magnetization component (red curve) and assumed modulation of
Ku (blue curves) as a function of t.
94
0.02
0.01CO
§ 0.00 N E
- 0.01
- 0.02
0 200 400Delay t (ps)
Figure 4.16 Measured Kerr signal change with 4 ps step size (black squares),
simulated polar magnetization component (red curve) and assumed modulation of
Ku (blue curves) as a function of t.
step size = 4 ps xr = 40 ps
305*
"rocg>co
a>
95
step size = 0.8 ps Tr = 2 ps ^0.001
ni 0.000
- 0.001
0 40Delay t (ps)
Figure 4.17 Measured Kerr signal change (black squares) and simulated polar
magnetization component (red curve) with rr = 2 ps.
96
Kerr
sign
al (
a.u.
)
0.02step size = 0.8 ps Tr = 15 ps / y *
0.01
* 0 . 0 0
- 0.01
0 40 80Delay t (ps)
Figure 4.18 Measured Kerr signal change (black squares) and simulated polar
magnetization component (red curve) with rr = 15 ps.
97
Kerr
sign
al (
a.u.
)
0.02
0.01CO^ 0.00
£
- 0.01
- 0.020 40 80
Delay t (ps)
Figure 4.19 Measured Kerr signal change (black squares) and simulated polar
magnetization component (red curve) with rr = 100 ps.
_ - step size = 0.8 ps — xr = 100 ps
98
Kerr
sign
al (
a.u.
)
Chapter 5
Spin precession excitation in manganite thin films
Pronounced spin precessions are observed in a geometry with
negligible canting of the magnetization in FM Lao.6 7Ca0 3 3 Mn0 3 thin films using
the TRMOKE technique. The precession amplitude monotonically decreases
with increasing field, indicating that the coherent spin rotation may be
triggered by a transient exchange field and not by demagnetization and/or
anisotropy field modulation. W e attribute the transient exchange field to
emergent AFM interactions due to charge transfer and modification of the
kinetic energy of eg electrons under optical excitation.
99
5.1 Introduction
Control of coherent spin rotation in ferromagnets has received intense
investigation in the past decade for fast magnetization switching in novel
magnetic recording and spintronic devices [91,117], Recently, optical
excitations of coherent spin precessions have been demonstrated in various
FM metals, providing a new approach for fast spin manipulation. The
precessions are triggered by effective field pulses generated via fast
demagnetization [118,119], or transient modulation of anisotropy constants
[93-95], The demagnetization occurs due to fast dissipation of spin angular
momentum via efficient spin-lattice thermalization [83], whereas the
anisotropy constant modulation may be caused by the alteration of spin-orbit
interactions due to hot electrons and electron phonon thermalization [93],
creation of itinerant carriers [94], or coherent acoustic phonons [95],
Despite the different excitation mechanisms, both scenarios require
canting of the magnetization, i.e., an angle between magnetization and the
external applied magnetic field to induce the coherent rotation. This geometry
is also essential for spin precession in exchange biased magnetic structures
[96], In such systems, the instant laser heating destroys the exchange
induced pinning of the FM spins at the interface, and thereby triggers the
rotation of the magnetization to the new direction. Djordjevic et al. recently
proposed a new excitation mechanism, where localized spin flip may decay
into short wavelength spin wave excitations and then transfer to low energy
spin wave modes [120], The canted geometry also may not be important for
100
ferrimagnetic materials in which the spin precession is excited via inverse
Faraday effect using circular polarized light [104,121].
In this Chapter, observations of pronounced spin precessions are
reported in a geometry with negligible canting of the magnetization in FM
La06 7Ca0.3 3MnO3 (LCMO) thin films on NdGa03 (NGO) substrates. Using
TRMOKE measurements, large spin precessions are observed over a wide
range of magnetic fields (0 -2 .5 T) applied along the in-plane easy axis. The
precession amplitude monotonically decreases with increasing field, in
contrast to the precession induced by the transient anisotropy field in the
state of canted magnetization. These observations indicate that the coherent
spin rotation may be triggered by a transient exchange interaction and not by
demagnetization and/or anisotropy field modulation. The picosecond
modulation of the exchange interaction results from the emergence of AFM
interaction caused by charge transfer and alteration of kinetic energy of eg
electrons in LCMO under optical excitation.
101
5.2 Samples and Experiments
The LCMO films are epitaxially grown by pulsed-laser deposition on
NGO or SrTi0 3 (STO) substrates. An excimer laser with a wavelength of 248
nm, an energy density of 2 J/cm2, and a repetition rate of 5 Hz is used. The
growth rate is around 0.04 nm/pulse, and the thickness of the film was
determined by the nominal value. The substrate temperature is around 780
°C and the oxygen pressure was about 500 mTorr during deposition. The
films are cooled down in about 400 Torr of oxygen at the rate of 20 °C/min
after deposition. Unstrained films were obtained by annealing the films in a
quartz tube at 900 °C with a flow of 50 standard cm3/min pure O2 for 24 h and
cooled at the rate of 10 °C/min to room temperature. The films on NGO
substrates exhibit an inplane uniaxial anisotropy of 0.2 T at thickness of 6 0 -
100 nm [122],
TRMOKE measurements are performed with a Ti:sapphire amplifier
laser system providing 150-fs pulses at 1-kHz repetition rate. W e use 800-nm
pump pulses with a fluence of 1 mJ/cm2 to excite the magnetization in our
samples, and the time evolution of the magnetization is monitored with 400-
nm probe pulses, where the Kerr rotation is at a maximum [123], with
fluencies on the order of 0.1 mJ/cm2. The transient Kerr rotation Ad is
measured in a cross polarization configuration. Rotation of the analyzer
through the extinction angle leads to the opposite sign of Ad. For transient
reflectivity ( AR ) measurements we use the same pump and probe
wavelengths as in TR-MOKE.
102
5.3 Coherent Spin Precession via Photoinduced Anti
ferromagnetic Interactions in Lao.67Ca0.33Mn03
Figure 5.1(a) shows the time evolution of Ai? from a 100 nm thick
LCMO film at 25 K. AR exhibits a fast response with a rise time comparable to
the laser pulse width and a slow response with decay time of 100 ps. The
latter component is associated with the slow spin-lattice relaxation [124].
The time evolution of Ad is shown in Fig. 5.1(b). Here, the external field
H is applied perpendicular to the sample plane and the probe beam is
incident at an angle of 5 degrees with respect to the sample normal,
permitting sensitive detection of the out-of-plane magnetization component
Mz. In contrast to the fast electronic response, the magnetic response shows
no instantaneous change on the time scale of 1 ps. The major feature of the
magnetic response manifests as strong oscillations which correspond to the
coherent precessions of the FM spins. W e therefore conclude that no
significant demagnetization occurs before the magnetization rotation is
launched. This finding is consistent with the observations of negligible
demagnetization in FM Lao.6Sr0.4Mn0 3 films at temperatures significantly
lower than its Curie temperature [123], The weak demagnetization is due to
the half metallic nature in manganites where the Elliot-Yafet scattering is
blocked [107],
103
3m
5
0 50 100 150
H I I Easy Axis (c)
Time (ps)
H x Plane (b)
3
■<3
0 200100
0.05 T
0.3 T
■ V
Time (ps)50 100 150 200
Time (ps)
3
% H = 0 T
0 200 400 600 0Time (ps) Frequency (GHz)
Figure 5.1 Transient reflectivity AR (a), and transient Kerr rotation Ad of LCMO film
with magnetic field H (0.5 T) applied perpendicular to the sample plane (b), with H
applied along the in-plane easy axis (c), and with zero field (d) at 20 K; Fourier
transform (e) of Fig. 5.1(d).
104
easy axis (b)
560 Oe
OOe
200 4000
hard axis (a)
« 560 Oe
200 4000Tim e (ps) Tim e (ps)
Figure 5.2 Transient Kerr rotations of Fe film with magnetic field applied along the in
plane hard axis (a) and easy axis (b) at RT.
105
Ferromagnetic metal Half metal
Figure 5.3 Three-temperature model in ferromagnetic metals and half-metals.
As a comparison, Fig. 5.2(a) shows the laser induced magnetization
dynamics in an Fe film at an external field of 550 Oe applied along the in
plane hard axis at room temperature (RT). The transient Kerr rotation clearly
shows a fast response which occurs within 1 ps. This response originates
from the fast demagnetization due to the spin flip process associated with the
spin-lattice interactions [118]. The demagnetization decays within 50 ps, while
the subsequent magnetization precessions persist for much longer time (>
500 ps) [125]. The excitation of the magnetization precessions in the Fe film
can be reasonably explained by the modulation of the anisotropy fields
caused by the demagnetization [118,119]. Under the laser interaction, the fast
demagnetization modifies the magnetocrystalline anisotropy, generating a
transient effective field Htr perpendicular to the static effective field Hen, and
thus the magnetization precessions are triggered.
Hence, the precession is significantly weaker and difficult to observe if
the external field is applied nearly close to the easy axes along which the
magnetization is pinned, as shown in Fig. 5.2(b). These observations clearly
indicate that magnetization canting and fast demagnetization are of crucial
importance in exciting spin precession in the Fe film, typical for 3d transition
ferromagnetic metals. Figure 5.3 depicts the difference of three-temperature
model in ferromagnetic metals and half-metals.
In contrast, Fig. 5.1(b) shows that fast and large demagnetization is not
present in the spin precession dynamics of LCMO films. In the following, we
will present further measurements of the magnetization dynamics when H is
107
applied along the uniaxial easy axis in the film plane, i.e., a configuration with
negligible canting of magnetization. The results are shown in Fig. 5.1(c).
Similar to the case where the magnetization is pulled out of plane, no fast
demagnetization occurs when spins are aligned within the film plane, while
pronounced spin precession is present over a wide range of magnetic fields.
Particularly, a large precession is also excited at zero field and persists for
over 1 ns, as shown in Fig. 5.1(d). We note beating patterns in the transient
Kerr spectra at zero field. Fourier transform reveals two spin waves modes at
frequencies of 11 GHz and 15 GHz (Fig. 5.1(e)). The mode with lower
frequency corresponds to the Kittel mode, as the detailed analysis of the
magnetic field dependence of the precession frequency shows [122]. The
higher-frequency mode may correspond to a surface Damon-Eshbach-type
mode or a higher-order standing spin wave mode within the ferromagnetic
film [126,127], The slow decay of the precession indicates that the dephasing
process is weak, and thus indicates a well-defined easy axis along which all
exchange-coupled spins are aligned at zero external field. In this case canting
of magnetization does not exist, and therefore modification of the in-plane
anisotropy may not be the main contributor to the transient field that launches
the spin precession shown in Fig. 5.1(c).
In the state of canted magnetization, anisotropy modulation will indeed
generate a transient field, triggering the spin precession. The anisotropy
modulation may arise from the thermal modification of magnetocrystalline
energy at elevated electron temperature [93]. Although the modulation of
108
anisotropy field Ha is independent of the applied field H, the transient torque
on the magnetization becomes stronger with increasing H as it pulls the
magnetization away from Ha, thus leading to larger precession angles at
higher fields. When the magnetization is nearly aligned along H at the field
strength stronger than Ha, the torque starts to saturate and the precession
angle decreases with increasing H. Thus, one can expect that the precession
amplitude reaches a maximum at the field approximately equal to the static
anisotropy field. The top left panel of Fig. 5.4(a) shows the field dependence
of precession amplitude obtained when H is applied at an angle of 45 degree
to the easy axis, a state with canted magnetization. We note that the
amplitude is sharply enhanced with increasing field and then slowly reduced
at fields larger than 0.3 T, consistent with our expectation. Similar results are
obtained from the other LCMO samples, as well as the Fe film with the field
applied along the hard axis and the direction at an angle of 10 degree to the
easiest axis (Fig. 5.4(b)). The modulation of the magnetic in-plane anisotropy
in an oxide ferromagnetic material C r02 also leads to similar dependence of
precession amplitude on H applied along the hard axis, and no precession is
detected if H is oriented close to the easy axis [93,128].
In contrast, the precession with the applied field along the easy axis of
LCMO shows the opposite dependence; i.e., the amplitude decreases with an
increasing field for all three LCMO samples, as shown in the right panels of
Fig. 5.4(a). This result confirms our previous conclusion that the transient
anisotropy field due to demagnetization or modification of anisotropy constant
109
is not the major contributor to launch the spin precession when H is applied
parallel to the easy axis. W e thus propose a transient exchange field H lrx that
triggers the spin precession.
(a)2 ■
«n3
. cSrEPn .D a hard axis■ □ □
□,S 60 nm
0.9
0.6
0.3
1-O
■%9 0o
LCMOeasy axis
□□
X1 i
1
1.2 o
D □ • □ 0.8 o
^ 100nm0.4
o o
1.2
■ ■ □ □
□ 0.8o
o. □g 100 nm* □...1.... ■.....A-... .......L
0.4J— ......
oOO
.1..... -..... 1...1 2 0
Magnetic Field (T)
(b>
T>
O.E<
1 •
□ □ 6 £0oo01
□□
cs
■'—r □
1□ 4 o Fe
-o.q Iff to easy axis1 . 1 i 1
hard axis□I * ....A - 1
2
0 400 800 0 400 800Magnetic Field (Oe)
Figure 5.4 Amplitude of precession of 60 nm and 100-nm LCMO films with field
applied at 45 degree to the easy axis and along the easy axis at 20 K (a), and Fe film
with field applied at 55 degree to the easy axis and of 10 degree to the easy axis at
RT (b).
110
In order to obtain the field dependence of H lrx , we simulate the
precession of the magnetization and follow its out-of-plane component. In the
simulation, the magnitude of //|£ is adjusted so that the calculated spin
precession angle coincides with the nominal precession angle obtained from
the measured precession amplitude (Fig. 5.4(a)). W e thus determine the field
dependence of the magnitude of H lrx for the 60-nm thick film, as shown in Fig.
5.5(a). We note that Hlrx is reduced with increasing applied field. This
dependency is a characteristic of the transient exchange field in manganites
as discussed below. In contrast, the transient anisotropy field H£r , which is
perpendicular to the magnetization, due to demagnetization or reduction of
the anisotropy constants is enhanced with increasing field. We present a
calculated result in Fig. 5.5(b), which corresponds to t f ' a in a configuration
with slight canting of magnetization, i.e., an angle of 5 degrees between the
applied field H and the uniaxial anisotropy field Hu, and 10% reduction of Hu
(0.2 T) after laser interaction.
^ 10
50o>
Z 40
30
0 1 2 M agnetic R eid (T )
0 1 2 M agnetic Field (T )
Figure 5.5 Transient exchange field (a) and transient anisotropy field (b) in a state of
canted magnetization (magnetic field at 5 degree to the easy axis) of 60-nm thick
LCMO film. The dashed line is a guide to the eye of the data.
I l l
The transient exchange field may result from the modification of the
exchange coupling because the optical excitation may change the balance
between the FM metallic and AFM insulating states in the film. The
coexistence of distinct FM metallic and AFM insulating phases in doped
manganites has long been established [129-131], and the AFM domain may
be melted in the presence of an external field. In the case of LCMO, several
reports indicated the presence of the AFM phase down to the lowest
temperatures [132-134], Very slow photoinduced drops in conductivity and
demagnetization have been observed in LCMO and interpreted as the
photoinduced shift in the FM/AFM phase balance triggered by the
photoexcited carriers that promote the cooperative Jahn-Teller distortions and
formation of AFM insulating clusters [135]. An ultrafast drop in conductivity
upon photoexcitation with femtosecond pulses was observed by Averitt et al.,
which is also consistent with the promotion of AFM correlations [124],
We speculate a formation of nanometer- to sub-micrometer- scale
regions with transient AFM interaction after the intense optical excitation,
similar to the size of AFM domains observed in quasistatic phase separated
manganites [129-132], Since the AFM regions are randomly distributed in
LCMO films, we would expect a spatial inhomogeneity of the transient
exchange field which may lead to excitations of the Damon Eshbach mode or
the standing spin waves [127,128]. Because application of the magnetic field
may suppress the AFM interaction, the magnitude of Hlrx is reduced at higher
fields as shown in Fig. 5.5(a).
112
The reorientation of the spins to form AFM clusters may take a long
time due to the slow spin-lattice thermalization process, a general feature in
manganites [136], to dissipate spin angular momentum. However, the
exchange interaction may be modified in a much faster time regime. This is
because the local carrier concentration and its kinetic energy play a crucial
role in determining the magnetic phases [137,138], Both photoinduced charge
transfer between neighboring Mn ions [139,140] and the thermalized
electrons may suppress the double exchange interaction and promote the
AFM coupling. Therefore, the exchange field pulse may be generated within
electron phonon relaxation time, much faster than the emergence of the
demagnetization caused by the real formation of AFM clusters.
This finding is corroborated by the temperature dependence of the
magnetization dynamics shown in Fig. 5.6. At temperature lower than 150 K,
the magnetization dynamics show negligible demagnetization. With
increasing temperature, a slow demagnetization (> 500 ps) is clearly
observable. However, the spin precession is launched nearly at the same
time at all temperatures lower than the Curie temperature Tc (~270 K). The
reduction of the precession amplitude at higher temperatures is mainly
caused by the decrease of the saturated magnetization. The reduction of the
magnetization also results in a pronounced decrease of precession frequency
at temperatures approaching to Tc, as shown in the inset of Fig. 5.6. This
result therefore supports our model that the precession is triggered by the
transient exchange field in the absence of spin-lattice thermalization.
113
■
3a
• w
<J
' m T=150 K
J j a -
• • • - • i -
T=230 K
r '.
|V T=260 K1
A
i ....». i ...... ..... —i—
■ -.2 0 N5 10
T=2
.............. 'SHMX5 T
) 100 200 3(T (K)
90 K
i...- ■.....
)0
10 250 500 0 250 500
Tim e (ps)
Figure 5.6 Transient Kerr rotations of LCMO film with applied field H=0.5T at various
temperatures. The inset shows the uniform precession frequency as a function of the
temperature.
114
A recent study reveals a transient anisotropy field due to the optical
induced transition between easy in-plane and easy out-of-plane anisotropy
states in La2-2xSri+2xMn20 7 bilayered manganites [94], However, we point out
that the LCMO film is not located at the boundary of two ferromagnetic
phases with different anisotropies, and therefore no transient field due to the
emerging of anisotropy field along different directions is expected to be
present in our LCMO films. Such a transient field also would be independent
on the external field, different from H lrx which we obtain from the simulation of
the field dependence of the precession amplitude as shown in Fig. 5.5(a).
115
Chapter 6
Magneto-electric coupling at PbZr0.52Ti o.4803/Lao.67Sro.33Mn03
interface
The interfacial spin state of the multiferroic heterostructure PbZr052 Ti
o.4803/Lao.67Sr0 3 3 Mn0 3 (PZT/LSMO) and its dependence on FE polarization is
investigated with MSHG technique at 78K. The spin alignment of Mn ions in
the first unit cell layer at the heterointerface can be tuned from FM to AFM
exchange coupled, while the bulk magnetization remains unchanged. Multiple
domains of both phases coexist as the ferroelectric polarization is switched.
The results will help promote the development of new interface-based
functionalities and device concepts.
116
6.1 Introduction
For the past few decades, magnetic recording has been a major
technology for information storage, while binary information is stored as
magnetization of magnetic elements. During this period of miniaturization of
magnetic memory cells, the reading procedure turns out to be scalable down
to the nanometer range. However, the traditional writing mechanism driven by
a magnetic field appears as a crucial limiting factor, due to the large power
consumption and loss of accuracy caused by the magnetic stray fields
affecting neighboring cells. In the modern MRAMs, spin-polarized current
induced spin-torque is used to switch magnetization states instead of
magnetic field, which somehow enables further down scaling of unit cells [9-
12]. Nevertheless, the large current densities required also lead to a
significant energy loss from heating. Thus in the push for low-energy
consumption logic devices, electric field control of magnetization switching via
strong magneto-electric coupling seems to offer a new mechanism to break
through those limitations and improve device performance; as such, it has
recently drawn increasing research interest.
Transition metal oxides are promising candidates for multifunctional
devices, due to their strong correlated electron system with coupled charge
[141], spin [142] and orbital [143] degrees of freedom. This route is carried
out by fabricating well-defined interfaces between these complex oxides, to
engineer and cross-couple their unique electric, magnetic, and transport
properties [144, 145]. The interfaces of TMO heterostructures have the
117
genuine property of breaking space inversion symmetry, and thus offer a
unique and important test-bed to promote new phases and properties that can
be controlled at atomic precision, such as conducting electron gases [141],
superconductivity [146] or electric polarization dependent spin transfer
[15,147],
The interfacial spin configuration of oxide heterostructures under
electronic and structural reconstruction is key to emerging multiferroic and
spintronic technologies with new functionality, and thus attracts research with
different approaches [148-152], However, most experimental probes are
indirect due to the lack of sufficient resolution to locate spin configuration at
the atomic scale. The interfacial spin state is often not clear, which prevents
further interpretation of the coupling between spin and other ordering
parameters at oxide heterointerfaces.
A good candidate for the magnetic constituent of such interface is
doped manganite, which received detailed understanding on carrier filling and
orbital effects [153, 154], To realize electronic and structural reconstructions
of doped manganite, electrostatic and strain effects are primary methods,
which modulate the competition between different interactions [155-160], In
this letter, the spin state at the PZT/LSMO interface is directly probed by the
interface-specific MSHG technique, and the control of exchange coupling
between interfacial spins is realized through their coupling with the charge
degree of freedom, which provides previously unknown critical insight into the
magnetic order at this multiferroic heterointerface.
118
6.2 Samples and experiments
The PZT/LSMO heterostructures used in this study are grown by
pulsed laser deposition on (100) LaAI03 (LAO) substrates. The LSMO layer
is grown at 600 °C under an oxygen pressure of 80 mTorr, using a laser
energy density of 1.8 J/cm2 and repetition rate of 10 Hz, followed by
annealing at 700 °C for 30 min in oxygen. The PZT layer is then deposited on
the LSMO layer at 600 °C at the same conditions. After deposition the
heterostructure is annealed at 700 °C for 30 min in oxygen at a pressure of
300 Torr. Finally, the films are cooled down to room temperature at a slow
rate. The film thickness of PZT and LSMO layer is around 550 nm and 100
nm respectively, determined from an X-P-200 profilometer. FE PZT layer is
used as a constituent of the heterostructure to adjust the carrier injection at
the interface layer, considering its high Curie temperature Tc [161], high
remnant polarization and low coercivity field £ c [162]. Such thick PZT layer is
used to suppress leakage current. The ferromagnetic LSMO layer is known to
exhibit high Curie temperature of about 370 K, colossal magneto-resistance
properties and half metallic behavior [155,163], and typically is used as a
spin-polarized detector [164], Indium tin oxide (ITO) top electrode is deposited
on top of the PZT layer to enable the application of an electrical field.
To investigate the interfacial magnetic state of the PZT/LSMO
heterostructures, we carry out MSHG measurements with fundamental laser
beam at 800-nm wavelength and detection of MSHG signal at 400-nm
wavelength, as depicted in Fig 6.1(a). The MSHG signal can be generated in
119
magnetic materials in which both space-inversion and time-reversal symmetry
are simultaneously broken [26-28, 165-167]. The broken time-reversal
symmetry results from the presence of magnetization, while the broken
space-inversion symmetry occurs only at the surface or interface of
centrosymmetric materials like LSMO. Considering these requirements, the
first unit cell layer of LSMO at PZT/LSMO boundary can generate MSHG
signal. However, the Nth (N>2) unit cell layer from such boundary will not
generate MSHG signal, as it is surrounded by the (N-1)th and (N+1)th unit cell
layers which preserve space inversion symmetry (Fig. 6.1(b)).
For comparison, MOKE measurements are carried out to probe the
magnetic state of LSMO bulk layer. MOKE technique is carried out with
similar experimental configuration as MSHG (Fig. 6.1(a)), except that the Kerr
signal (800 nm) is detected with a high-speed silicon detector after a P-
polarized analyzer. The optical absorption coefficient in LSMO for SHG light
(A = 800nm) is about 1 * 105cm-1 . Considering the exponential decay of light
intensity, the magnetic contribution of the first three unit cell layers to the
MOKE signal which probes the magnetization of the entire LSMO film is less
than 2%, which is close to the noise level of the MOKE hysteresis loop.
Hence, the MOKE technique is not sensitive at all to the change of interface
magnetization.The external magnetic field H is applied along the [100] axis of
LSMO layer for both MOKE and MSHG.
Figure 6.2(a) shows the X-Ray Diffraction results of PZT/LSMO
heterostructure measured with a Siemens D500 x-ray diffractometer (Cu Ka
120
radiation) in a 0-20 scan. Only the (100) diffraction peaks of LSMO and PZT
are observed, along with those of the single crystal LaAI03 substrate. This
suggests that the individual ferromagnetic and ferroelectric layers are grown
epitaxially block-by-block on the substrate and form smooth films without any
secondary parasitic phases. The LSMO layer is under slight tensile strain with
lattice parameter a=3.885 A0, due to the oxygen pressure during growth [168,
169]. Figures 6.2(b)-6.2(d) show RHEED patterns of LAO substrate and
LAO/LSMO structures before and after the deposition of PZT layer on top.
Very intense specular spots, Kikuchi lines and Laue zones are observed,
revealing that the LAO substrates used in this work are high-quality single
crystals and perfect in-plane epitaxial growth of the LAO, LSMO and PZT
layer with smooth surface.
121
(a) MSHG MOKE
PZT
Interface
LSMO
• Ti Zr
PZT
Interface
LSMO
Figure 6.1 Schematic of MSHG and MOKE measurements (a), and depiction of
PZT/LSMO interface on atomic scale (b). MSHG probing plane is marked with purple
color.
122
oo
CM
20 30 40 50Braggs Angie (2e)
(b) (c) (d)
LAO RHEED pattern ISMO/LAO RHEED pattern PZT/LSMO/LAO RHEED pattern
Figure 6.2 XRD results of PZT/LSMO heterostructure (a), RHEED patterns of LAO
substrate surface (b), LSMO surface before deposition of PZT layer on top (c) and
PZT surface of PZT/LSMO/LAO heterostructure (d).
123
6.3 Charge control of antiferromagnetism at PbZr0 .s2Ti 0 .4 8 O3/
Lao.6 7Sr0.33Mn0 3 interface
Figures 6.3(a) and 6.3(b) display the interfacial and bulk magnetization
as a function of applied gate voltage Ug probed with MSHG and MOKE at 78K,
respectively. The key finding is that the magnetization of the interfacial layer
is modulated by Ug, while the bulk magnetization is not. The magnetic
contrast for MSHG and MOKE is defined as
„ _ /(+M )-/(-M ) /cA ~ r(+M)+n - m ( 6 1 )
where l(±M) is the signal intensity received by the detector when the
macroscopic magnetization M is aligned with external magnetic field in either
positive or negative [100] direction (Fig. 6.4(a)). The magnetic contrasts
obtained by both MSHG and MOKE is linear with M in the probed volume, if
magnetic component is much smaller than non-magnetic component in signal
intensity as in our case [26], The magnetic contrast A of MSHG measurement
as a function of Ug is displayed in Fig. 6.4(b), where positive Ug refers to
positive voltage applied to the ITO top electrode. These results indicate that
the modulation of gate voltage on PZT/LSMO heterostructure, which is not
strong enough to change the magnetization of the bulk LSMO layer, can
significantly tune the magnetic state at the interface. This may provide a clue
for fabricating low energy consumption multifunctional devices based on
similar material systems.
124
interface10 10
0 P o
7.5 V -7.5 V
-10 V -10 V- 10-10
-1000 1000 -1000 1000Magnetic field (Oe) Magnetic field (Oe)
Figure 6.3 Interface (a) and Bulk (b) magnetic hysteresis loops from PZT/LSMO
heterostructure for different external gate voltages Ug measured with MSHG and
MOKE techniques at 78 K, respectively.
125
MOKE
sig
nal (
a.u)
0.36
3.0.32
0.28
0.1
<
0.0
Figure 6.4 MSHG hysteresis loop with Ug = +10V indicating l(-M) and l(+M) used to
determine the magnetic contrasts from Eq. (6.1) (a), and MSHG magnetic contrast
A as a function of Ug (b). Increasing (decreasing) gate voltages, Ug up (Ug down), are
labeled in red (black). The direction of positive Ug is defined in the inset.
Ug down Ug up
Ug (V)
I (+M)
l(-M)
1000-1000 0H (Oe)
126
(a)
Ug down Ug up
-60
-15 0 15Ug(V)
(b)
Ug down Ug up0.1
0.0
60-60 0
P (jaC/cm2)Figure 6.5 Ferroelectric polarization (P) in PZT as a function of Ug (a). The direction
of positive P is defined as pointing towards LSMO thin film. M SHG magnetic contrast
A as a function of P of PZT (b). The dashed line in Fig. 6.5(b) is the least square fit
for the data points.
127
(a)
OOo
CO
4
-10
(b)
0ug(V)
10
Ug down Ug up
- + — --
0.0
3 4S (1/1000)
Figure 6.6 Strain state (S) of LSMO interface as a function of gate voltage (a). MSHG
magnetic contrast A as a function of strain (5) of PZT (b).
128
To elucidate the change of interface magnetization with gate voltage,
FE polarization (P) measurements are carried out, as shown in Fig. 6.5(a).
The FE polarization saturates with gate voltage above 10V (electric field of
approximately 180 kV/cm) or below -10V, indicating that all FE domains are
aligned. The switching transition around Ug = ±2V is not sharp, resulting from
the variance in switching threshold of different FE domains [170]. The hole
density at the PZT/LSMO interface can be electrostatically modulated, i.e.
hole depletion/accumulation occurs when P points towards/away from the
LSMO layer [149-151]. From the FE polarization measurement, we derive the
change in magnetic contrast A as a function of P, as shown in Fig. 6.5(b),
using the dependence of A on gate voltage Ug (Fig. 6.4(b)). The relationship
between A and P exhibits a linear correlation, independently of the history of
polarizing PZT layer, as the data with decreasing gate voltage (black)
overlaps with the increasing one (red).
The lattice constant of PZT layer can also be simultaneously
modulated by gate voltage application, due to the piezoelectricity of PZT. To
estimate the influence of gate voltage (Polarization) on interface strain state,
we used the lattice distortion of PZT layer from Ref. [171]. In this estimation,
we assumed that the modulated x-y plane strain of PZT by gate voltage is
transferred onto LSMO lattice at interface. This assumption might lead to a
larger estimation of LSMO lattice distortion, but is still similar with realistic
case. The strain S is defined as the percentage variation in the in-plane
lattice parameter of LSMO layer relative to the bulk LSMO value, and
129
displayed in Fig. 6.6(a) as a function of gate voltage. The offset at zero
voltage (about 0.3%) is based on the lattice constant derived from XRD
measurement. The strain S at interface is weakly modulated by gate voltage
Ug, and its dependence on gate voltage Ug exhibits a butterfly shape.
Therefore, the magnetic contrast A as a function of strain S is depicted in Fig.
6.6(b), which is derived from the relationship between strain and gate voltage
in PZT thin films studied in Ref. [171] and using the dependence of A on gate
voltage Ug (Fig. 6.4(b)). There is no clear correlation between interface
magnetization and strain.
Thus, the observed change of interface magnetization with gate
voltage results from charge injection. The charge mediated mechanism is
depicted by a schematic model of screening charge accumulation at the
PZT/LSMO interface (Fig. 6.7), responding to opposite FE polarization
orientations in the PZT layer. The polarization follows the relative
displacement along the z direction of the metal cations (Pb, Zr and Ti) with
respect to their neighboring oxygen anions in the same (001) plane [151].
When the FE polarization of PZT is pointing towards the LSMO layer,
electrons will accumulate at the LSMO surface to compensate for the charge
on the PZT side. On the other hand, when the FE polarization of PZT is
pointing away from the LSMO layer, holes accumulate at the surface of
LSMO. To quantitatively determine the hole doping level p modulated by gate
voltage at the first unit cell layer of interface LSMO, we assume that the
Polarization (P) of PZT initiates the charge at geometric boundary between
130
PZT and LSMO (z=0) as P 5 (z ) , and the density of screening charges
accumulating at the LSMO side of interface exponentially decays with the
distance away from PZT/LSMO geometric boundary, as depicted in Fig. 3.
Thus the charge distribution can be written as P8(z) e~z/x a t z > 0 ,A.
which is at the LSMO side of interface. The Thomas-Fermi screening length
X = yje/NEfe2 = 4.14 >4°, where s = yj32e0 is the dielectric constant, NEf =
0.6(ev~1u.c.~1) is the Fermi level density of states [172]. The modulation of
hole doping level per Mn ion at single-unit-cell interface LSMO layer is
approximately calculated from the integration of charge density from z=0+ to
z= lattice constant of LSMO along [001] direction, and the degree of which is
calculated to be ±0.231 electrons per Mn ion in the first unit cell layer at the
interface corresponding to ± 10V gate voltage.
PZT(z>0)
e l i p r - @y LSMO/ (? < 0) 1
(z = 0)
Figure 6.7 schematic model of ions displacement and screen charge accumulation
under opposite polarization state.
The loss of interface magnetic contrast A by hole injection (negative Ug)
can be due to either a modulation of saturation magnetization, or changes in
the coupling mechanism between Mn ions, or changes in the
magnetocrystalline surface anisotropy, or a combination thereof. The spin
alignment of Mn 3d electrons still obeys Hund’s rule, and the magnetic
moment per Mn site cannot be smaller than 3 Bohr magnetons, as the charge
injection by PZT can only modulate the eg electron concentration [149, 172],
Thus, the depletion of electrons with coherent spins by hole injection can
cause a reduction of magnetization, but the degree of which is limited and
thus not responsible for the significant change of interface magnetic contrast
by one order of magnitude.
Another possibility is that the surface magnetic anisotropy changes
with increasing hole doping and the magnetization starts to point out of plane.
In the probe with optical polarization configuration S in S out, the magnetic
tensor X y y y ( M x ) is only sensitive to the Magnetization changes in the
longitudinal direction, and can not exclude such possibility. In addition, we
carried out MSHG measurement with P in S out optical polarization
combination, where magnetic tensor contains both XyzziMx) . Xyxx{Mx)
sensitive to longitudinal magnetic moment and Xyzx(M z) sensitive to polar
magnetic moment. The results under opposite polarization states of PZT are
displayed in Fig. 6.8. Under -10V gate voltage, there is no magnetic contrast,
hinting not only M x but also M z become zero. There is still a possibility that
surface magnetic anisotropy change as hole doping is modulated. But in our
132
case, we pull the interface macroscopic magnetization aligned with strong
enough magnetic field. This speculation is also supported by the fact that the
magnetization in magnetic hysteresis loop is saturated above 300 Oe. Thus,
we conclude that the interface magnetization is reduced to zero when
negative gate voltage is applied.
t 136 ’</> tz0)
■M
i 1.32cg>'</)<D S 128
-600 0 600 H (Oe)
Figure 6 .8 P in S out M SHG results under different polarization states.
Therefore, the only reason is a change of exchange coupling at the
PZT/LSMO interface by hole injection as pointed out by Vaz et al. [149]. The
diminishing magnetic contrast in MSHG measurement with increasing hole
doping suggests the appearance of AFM spin alignment in the first unit cell
layer at the PZT/LSMO interface.
—- +10 V U ■ "1"1 * ■— -10V 1 k
133
Next, we discuss the evolution of interface spin alignment in terms of
3d orbital occupancy of Mn ions. The origin of the AFM phase at the interface
is not completely understood. A recent discovery shows that d3z2 _ r 2 orbital
has lower energy than d x2 _y2 orbital, and thus is more occupied with electrons
at tensile strained LSMO surface or interface as a result of broken space
inversion symmetry [173-176], in contrast to tensile strained LSMO bulk.
According to Ref. [173], the d x2 _ y2 orbitals are unoccupied at the interface,
which results in anisotropy of exchange coupling [148]. Double Exchange
induced by hopping of eg electrons is stronger in the z direction (surface
normal), while super exchange induced by local t2 9 electrons is stronger in the
x-y plane (surface plane). It is demonstrated by C. Aruta et al. [175] that this
interfacial d 3z2 _ r 2 orbital occupation favors the C-type AFM spin ordering, and
the easy axis of the AFM phase is oriented along z direction, irrespective of
the strain conditions. Moreover, A. Tebano et al. [173] provided microscopic
evidence that this preference in orbital occupancy at the surface of LSMO
exists, independent of the chemical nature of the substrate and the presence
of capping layer. This is consistent with our observation that spin coupling on
x-y plane is tuned into AFM type at the PZT/LSMO interface, as the d x2 _y2
orbital becomes depleted by hole injection. Furthermore, as electrons
accumulate at the LSMO interface layer by tuning the FE polarization of the
PZT layer, in-plane orbital occupation of eg electrons is also expected,
corresponding to the appearance of FM coupling on the x-y plane. Further
injection of electrons into interface layer might lead to expansion of A-type
134
AFM regions at the expenses of the FM double exchange regions, which
resembles the AFM coupling of LaMn03[169]. Further studies are required;
particularly in view of the fact that spin transport properties may depend upon
the phase transition by charge injection or oxygen vacancy concentration at
the interface.
The transition from AFM to FM spin alignment at the PZT/LSMO
interface with increasing P is gradual (Fig. 6.5(b)), indicating that multiple
domains of both phases coexist as the ferroelectric polarization is switched.
Multiple domains with different polarizations exist near the gate voltage where
FE switching occurs in PZT thin film, as reported in Ref. [170]. Some FE
domains may be switched to the opposite polarization, which leads to a
distribution of local FM and AFM domains in the LSMO interface layer. When
P is saturated and points towards (away from) the LSMO film, the interface
spin alignment is FM (AFM) coupled. Thus, the redistribution of local
magnetic domains at the PZT/LSMO interface leads to a gradual change of
magnetic contrast in MSHG measurement.
135
Chapter 7
Conclusions
Fundamental understandings of magnetic properties and spin
dynamics help facilitate the design and fabrication of new magnetic devices
with desirable performance. In this dissertation, intrinsic Gilbert damping a0
and perpendicular magnetic anisotropy (PMA) in L1oFePd(i.X)Ptx ternary alloy
films can be continuously tuned by spin-orbit coupling strength f with
appropriate Pt/Pd concentration. In particular, aro and PMA are found to be
proportional to <f2 from TRMOKE studies for the first time, which is consistent
with ab-initio density functional calculations. The tunability of PMA is
enhanced by increasing chemical order in FePdi.xPtx alloys and by lowering
the temperature. Moreover, control of the anti-site disorder c with proper
growth temperature results in significant variation of aro in FePt thin films. As c
increases from 3% to 16%, aro increases by more than a factor of three from
0.06 to 0.19. The variation of c mainly affects the electron scattering rate 1/re,
while other leading parameters remain unchanged. A linear correlation
between croand 1/re is experimentally observed, consistent with the spin-orbit
coupling torque correlation model for interband transitions. The ar0 remains
unchanged with temperature, indicating that the intraband contribution to aro is
inundated by impurity scattering. Furthermore, as anti-site occupation
decreases, the perpendicular magnetic anisotropy increases and the Landau
136
g factor exhibits a negative shift due to an increase in orbital momentum
anisotropy.
The efficiency of spin precession excitation by laser pulses is
significantly improved through inserting an AFM CoO layer into Fe/MgO
heterostructure which establishes the uniaxial magnetic anisotropy Ku along
Fe [100] direction. The modulation of Ku by laser pump pulses generates a
fast exchange torque to the FM Fe magnetization. The transient torque is
enhanced at temperatures where Ku varies significantly and with external
magnetic fields competitive to the magnetic anisotropy fields, which results in
spin precession amplitude peaks. The excitation approach by modulation of
FM-AFM exchange interaction is much more efficient with 400-nm-
wavelength pulses via photo-excited charge transfer processes in AFM CoO
layer than modulating the FM order of Fe with 800-nm-wavelength pulses.
The recovery time of the exchange torque is around 40 ps, much faster than
the cooling time from demagnetization. Our results will help promote the
development of low energy consumption magnetic device concepts for fast
spin manipulation at room temperature.
This dissertation also presents photo-induced magnetization
precession in half-metallic Lao6 7 Cao3 3Mn0 3 thin films grown on NdGa0 3
substrates with TRMOKE technique. For the first time, pronounced spin
precessions are observed in a geometry with negligible canting of the
magnetization with a wide range of magnetic fields (0-2.5 T) applied along
the in-plane easy axis. The precession amplitude monotonically decreases
137
with increasing field, in contrast to the precession induced by the transient
anisotropy field in the state of canted magnetization. These observations
indicate that the coherent spin rotation may be triggered by a transient
exchange interaction and not by demagnetization and/or anisotropy field
modulation. The picosecond modulation of the exchange interaction results
from the emergence of AFM interaction caused by charge transfer and
alteration of kinetic energy of eg electrons in LCMO under optical excitation.
Finally, the interfacial spin state of PZT/LSMO heterostructure and its
dependence on FE polarization is investigated with magnetic second-
harmonic generation at 78K. The magnetization in the first unit cell layer at
the interface can be widely tuned over an order of magnitude by varying the
applied electric field between ± 1 8 0 kV/cm. The cross-coupling between
ferroelectric and ferromagnetic behavior at the heterointerface is mainly
electronic, due to charge injection. The ferroelectric polarization of PZT can
tune the spin alignment of Mn ions at the heterointerface from FM to AFM
exchange coupled. In contrast, the bulk magnetization probed with MOKE
remains unchanged. Multiple domains of FM and AFM phase coexist as the
ferroelectric polarization of PZT is switched.
138
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VITA
Xin Ma was born in Fuzhou, Jiangxi, China on May 18, 1988. Xin Ma
received his B.S. degree majored in Condensed Matter Physics at the
University of Science and Technology of China in 2009.
In August 2009, Xin Ma entered the College of William and Mary as a
research assistant in the Department of Applied Science. As a PhD candidate,
Xin Ma devoted to investigating magnetic properties and materials, including
spin dynamics and interface magnetization, with pulsed lasers and ultrafast
techniques.
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