1

Voltage Control of Magnetism

A Dissertation Presented

by

Ziyao Zhou

to

The Department of Electrical and Computer Engineering

in partial fulfillment of the requirements

for the degree of Doctor of Philosophy

in Electrical Engineering

Reviewer 1:Professor Nianxiang Sun

Reviewer 2:Professor Philip Serafim

Reviewer 3:Professor Yongmin Liu

Northeastern University

Boston, Massachusetts

April 2014

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Abstract

In past decades, attracted by the increasing demand of compact, fast, and low

energy consumption RF/microwave devices, many researchers have devoted their

efforts to realizing electric field control of magnetism, instead of magnetic field. For

instance, within traditional RF/microwave devices, ferromagnetic resonance are

controlled by bulky, noisy, slow and energy consumption electromagnets. This limits

its application in many important, low mass and energy consuming requirement

carriers, such as aircraft, satellites, radars and communication devices. As a result,

novel functional material, which can be integrated into non-volatile, light, and

energy-efficient electronic devices, need to be discovered. Multiferroics, a composite

material combined with ferromagnetic material and ferroelectric material, is widely

studied as a great candidate for E-field tunable RF/microwave applications like

tunable resonators, phase shifters, tunable inductors and tunable filters. The

coexistence of ferroelectricity and ferromagnetism in multiferroics introduces

interaction between ferroelectric property and ferromagnetic properties, therefore,

allowing electric field (E-field) control of ferromagnetism through varying

mechanism. In our work, different mechanism-based magnetoelectric (ME) coupling

in multiferroics heterostructure was investigated for the development of novel

generation, voltage-controllable, high-speed, compact RF/microwave devices with

greater energy efficiency.

Firstly, ME coupling was realized in different magnetic thin film/ferroelectric slab

heterostructures. By decreasing the saturation magnetization of Cr doping Ni

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magnetic thin film, large ME coupling in NiCr/PbZr0.52Ti0.48O3 (PZT) and

NiCr/PbZn1/3Nb2/3O2.4(PbTiO3)0.6 (PZNPT) was obtained. Furthermore, non-volatile

voltage impulse tunability was discovered through electric field-induced phase

transition in FeGaB/PZNPT multiferroics heterostructure. Giant ME coupling

coefficient ~3000 Oe cm/kV was observed at PZNPT phase transition points. In

FeGaB/Pb0.8Sn0.2Zr0.52Ti0.48O3 (PSZT) magnetic/antiferroelectric multiferroic

heterostructure, antiferroelectric-ferroelectric phase transition in PSZT substrate gives

us another opportunity to realize the voltage impulse tunable magnetic properties. The

non-volatile tunability with large ME coupling effect offers a great opportunity of

E-field control of magnetism in real RF/microwave applications.

Secondly, traditional deposition methods like sputtering, Pulsed laser deposition

(PLD), or Molecular beam epitaxy (MBE) require a high fabrication temperature

(>600 oC), which limits their application in integrated circuits. We used low

temperature(

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other mechanisms-induced ME coupling were also studied in our experiment. Large

interfacial charge mediated ME coupling effective field of 40 Oe was achieved in

Co0.3Fe0.7/Ba0.6Sr0.4TiO3 multiferroic heterostructure. The charge effect amplitude

dependence of magnetic film thickness was systematically investigated in

NiFe/SrTiO3 multiferroic heterostructure. Lastly, the ME coupling in CoFe/BiFeO3

(BFO) heterostructure induced by interfacial exchange coupling between CoFe

moment and canted moment in BFO was studied quantitively by FMR measurements.

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Acknowledgements

I would like to extend my greatest gratitude to my advisor Professor Nian X. Sun

for his guidance, encouragement and total commitment throughout my graduate career,

and also for the level of trust he demonstrated towards me. I would also like to thank

my dissertation committee, Professor Philip E. Serafim and Professor Yongmin Liu

for their suggestions and support.

Additional thanks are owed to all of my colleagues: Dr. Ming Liu, Dr. Yunume

Obi, Dr. Jing Lou, Dr. Xing Xing, Dr. Xi Yang, Dr. Ming Li, Dr. Shawn Beguhn, Dr.

Dazhi Sun, Dr. Zhongqiang Hu, Dr. Satoru Emori, Scott Rand, Yuan Gao, Tianxiang

Nan, Xinjun Wang for their help, time and wonderful discussions.

I am also extremely to Professor Ramamoorthy Ramesh from UC Berkeley and

his students Dr. Morgan Trassin, Gao Ya, Deyang Chen for their contributions. Also,

thanks to Prof. Carmine Victtoria, Prof. Vincent Harris, Dr. Yajie Chen, Dr. Bolin Hu,

Dr. Gail Brown, Dr. Brandon Howe, Dr. J. H. Jones and Dr. Krishnamurthy

Mahalingam from Air Force Research Lab and S. R. Bowden, D. T. Pierce and John

Unguris, from National Isititution of Standard and Technology, for their help with our

sample fabrications and measurements.

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Table of contents Abstract ......................................................................................................................... 2

Acknowledgements ...................................................................................................... 5

Chapter 1 Introduction of voltage control of magnetism in multiferroics

heterostructure ........................................................................................................... 11

1.1 Multiferroics and mangetoelectric materials ............................................. 11

1.2 Strain/Stress mediated ME coupling ........................................................... 13

Chapter 2 Magnetic/Ferroelectric multiferroics for tunable microwave

applications ................................................................................................................. 26

2.2 Non volatile tunable FeGaB/PSZT magnetic/antiferroelectric

heterostructures .................................................................................................. 36

2.2.1 FeGaB/PSZT multiferroic heterostructure fabrication .................. 36

2.2.2 Non-volatile control of magnetism in FeGaB/PSZT multiferroic

heterostructure ............................................................................................. 40

2.3 Non volatile tunable FeGaB/PZNPT magnetic/ferroelectric

heterostructures with giant tunability ............................................................... 44

2.3.1 FeGaB/PZNPT multiferroic heterostructure characteration ......... 44

2.3.2 RF/microwave tunability of FeGaB/PZNPT heterostructure ......... 48

2.3.3 Non-volatile switch of magnetism in FeGaB/PZNPT

heterostructure ............................................................................................. 53

Chapter 3 Low temperature fabricated multiferroics heterostructure ................ 58

3.1.1 ZnO and Al-doped ZnO thin film fabrication .................................. 60

3.1.2 ZnO and Al-doped ZnO thin film characterization ......................... 68

Chapter 4 Interfacial mediated magnetoelectric coupling in heterostructure

multiferroics ............................................................................................................... 89

4.1 Charge mediated ME coupling in NiFe/STO multiferroic heterostructure89

4.1.1 Thickness dependence of NiFe magnetic layer on STO layer ......... 90

4.1.2 ME coupling strength study on different NiFe/STO

heterostructures ........................................................................................... 93

4.1.3 Explaination of thickness dependence of ME coupling strength .... 98

4.3 Interfacial exchange coupling in CoFe/BiFeO3 multiferroic

heterostructure .................................................................................................. 102

4.2.1 CoFe/BFO multiferroic heterostructure fabrication and domain

pattern images ............................................................................................ 103

4.2.2 Eletric field induced ME coupling in CoFe/BFO ........................... 106

4.3.3 Modeling of canted moment in BFO switched by E-field ............. 109

Chapter 5 Conclusion and future work ................................................................. 118

5.1 Summary ...................................................................................................... 118

5.2 Further Research ........................................................................................ 120

References ................................................................................................................. 121

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List of Figures

Figure 1.1 The relationship between multiferroic and magnetoelectric materials.......14

Figure 1.2 Normalized Kerr rotation hysteresis curves (M-H) loops...........................16

Figure 1.3 M-H loops and FMR spectra of FeCoB/PMN-PT (011).............................18

Figure 1.4. Schematics of domain structures and reciprocal space maps (RSMs) about

(022) and (002) reflections of PMN-PT(011) under various applied electric fields....19

Figure 1.5 Electric-field-induced switching in CoFeB/MgO/CoFeB junction............24

Figure 1.6 Macro-spin model simulation of coherent magnetization switching under

various pulse duration conditions.................................................................................25

Figure 1.7 M-H loops and domain pattern of CoFe/BFO heterostructure...................26

Figure 1.8 Magnetoresistance measurements of CoFe/BFO under varied E-field.......27

Figure 2.1 M-H loops of NiCr thin film with different content...................................30

Figure 2.2 FMR spectra of NiCr thin film with different content................................31

Figure 2.3 FMR fields dependence of E-fields in NiCr/PZT.......................................33

Figure 2.4 FMR fields dependence of E-fields in NiCr/PZNPT..................................34

Figure 2.5 E-field dependence of NiCr Gilbert damping constant...............................36

Figure 2.6 Properties of FeGaB film deposited on the top or on the side of

Pb(Sn,Zr,Ti)O3 ceramics..............................................................................................40

Figure 2.7 M-H loops and FMR spectra under varying E-field of FeGaB/PSZT

multiferroics heterostructure........................................................................................42

Figure 2.8 Magnetization and FMR field switches of FeGaB/PSZT by voltage

impulses........................................................................................................................45

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Figure 2.9 Characterization of amorphous FeGaB on (011) orientated PZN-PT

substrate........................................................................................................................48

Figure 2.10 E-field tuning FMR properties of FeGaB/PZNPT....................................50

Figure 2.11 E-field induced FMR frequency shift under various magnetic fields.......52

Figure 2.12. Theoretical simulation (solid line) and experiment results (symbol) of

electric-field-induced FMR change..............................................................................55

Figure 2.13 Hysteresis loops of E-field vs. FMR frequency........................................57

Figure 2.14 E-field induced non-volatile switch in FeGaB/PZNPT............................58

Figure 3.1 SEM images of the ZnO microstructures...................................................64

Figure 3.2 Surface SEM images of the ZnO microstructures......................................66

Figure 3.3 SEM images of the ZnO microstructures with varying precursor conc......68

Figure 3.4 XRD patterns of the ZnO films...................................................................69

Figure 3.5 Optical absorption wavelength spectrum of the ZnO films........................71

Figure 3.6 XRD patterns of the Zn1-xAlxO thin films with varying Al concentration

(for x=0.02, x=0.06).....................................................................................................73

Figure 3.7 Typical XPS data of O1s in Zn1-xAlxO thin films and its Gaussian-resolved

component for x=0.06 Al concentration and x=0.02 Al concentration........................73

Figure 3.8 Plot of resistivity, hall mobility and carrier concentration as a function of

Al concentration (for x=0.002 to 0.02) for the Zn1-xAlxO thin films...........................74

Figure 3.9 Optical absorption spectra of ZnO with varying Al concentration.............76

Figure 3.10 Linear analysis confirmation between band gap energy and carrier

concentration................................................................................................................77

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Figure 3. 11 (a) X-ray diffraction pattern of the spin-spray deposited Fe3O4/ZnO thin

films multiferroics composite; (b) Energy Filtered TEM image showing the zinc oxide

(red) and iron oxide (green) layers; (c) HRTEM image of the iron oxide/zinc oxide

interface........................................................................................................................81

Figure 3.12. (a) The measured X-ray diffraction image using high-energy X-rays in

transmission geometry; (b) a 2 section from 2.1 to 6.2 over the azimuthal angle

from 0 to 180.............................................................................................................83

Figure 3.13. (a) Typical magnetic hysteresis loops of a spin-spray deposited

Fe3O4/ZnO ferrite/piezoelectric multiferroics heterostructure on glass substrate; (b)

In-plane magnetic hysteresis loops of the Fe3O4/ZnO multiferroics heterostructure

under different external electric voltages measured by VSM. The enlarged ME

coupling hysteresis loop shift is shown on upper left coordinate system; (c)

Out-of-plane magnetic hysteresis loops of the Fe3O4/ZnO multiferroics

heterostructure under different external electric voltages............................................84

Figure 3.14 (a) Piezoelectric coefficient measurements of ZnO thin film by Bending

Cantilever Beam Method; (b) Electric field dependence of the in-plane field-sweep

FMR spectra of the Fe3O4/ZnO multiferroics heterostructure measured at 9.3 GHz.

The zero cross part was enlarged to demonstrate a clear ME coupling shift at bottom

right inset; (c) X-band in-plane ferromagnetic resonance (FMR) field of the

Fe3O4/ZnO multiferroics heterostructure measured at varying applied voltages across

the ZnO layer................................................................................................................88

Figure 4.1 (a) X-ray diffraction pattern of the RF sputtered STO/Pt multilayer on Si

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substrate; (b) Polarization vs. Electric field of STO thin film; (c) AFM image of STO

surface with calibrated roughness of 0.88 nm..............................................................94

Figure. 4.2 Schematic of the sample used for a voltage-induced FMR field change in

Cu/NiFe/STO/Pt/Si.......................................................................................................95

Figure 4.3 (a) FMR field and calculated perpendicular energy dependence of NiFe

thickness; (b) FMR field shift dependence of applied voltage at varying thickness of

NiFe thin film on STO layer. (c) FMR effective field shift vs inverse of thickness

under varying voltage gaps...........................................................................................99

Figure 4.4 (a) E-field induced magnetic surface anisotropy δKS change and

normalized magnetic surface anisotropy per effective surface area δKS/Seffect

dependence of applied voltage bias; (b) Schematic of NiFe thin film deposition

dynamics surface at varying thickness.......................................................................100

Figure 4.5 (a) A schematic of the CoFe/BFO patterned multiferroic heterostructure.

The E-field was applied perpendicular to the CoFe disk. (b) The ferroelectric domain

structure imaged using the BSE intensity. (c) The simultaneously acquired SEMPA

image of the magnetic structure. The magnetization direction, θxy, is represented by

color as indicated by the color wheel. (d) An enlarged view of the magnetization in (c)

with arrows added to show the measured magnetization directions and the relationship

to the DSO substrate...................................................................................................105

Figure 4.6 FMR fields measurements of CoFe/BFO multiferroic heterostructure....107

Figure 4.7 Modeling and FMR spectra of CoFe/BFO multiferroic heterostructure...110

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Chapter 1 Introduction of Voltage Vontrol of Magnetism in

Multiferroics Heterostructure

Voltage control of magnetism is an extremely important technology in magnetic

mediated data storage, sensors and spintronics device. Traditional RF/microwave

devices were tuned through a current generated magnetic field. This tuning

mechanism is noisy, bulky, slow, and energy inefficient, which limits its application in

real devices.

Multiferroics, a combination of ferromagnetism and ferroelectricity properties,

which are coupling to each other, have created a lot of novel physical phenomena. The

coexistence of magnetization and polarization in multiferroics allows the electric field

control of magnetism or magnetic field control of polarization. This new mechanism

could lead to a new generation of memory devices, with four-state logic in a single

device, high frequency microwave and spintronics devices due to E-field controllable

of magnetism in these materials [1, 2, 3, 4, 5, 6]. In the first chapter, we will introduce

the basic principle of E-field control of magnetism in multiferroics, including

previous works in multiferroic heterostructures and the correlated potential

applications are also discussed.

1.1 Multiferroics and mangetoelectric materials

In the past decades, multiferroics have attracted many interests due to its significant

improvement to data storage, sensors and spintronics[7, 8, 9] devices. In the definition

of multiferroics, ferroelectricity is a spontaneous electric polarization in certain

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material that can be controlled by applying an external electric field; Ferromagnetism

is a spontaneous magnetic polarization in a material that can be controlled by

applying an external magnetic field. Fig 1.1 shows the relationship between

multiferroic and magnetoelectric materials [2]. The red area shows multiferroic

material, consisting with ferroelectric and ferromagnetic properties. The

cross-coupling, so called magnetoelectric (ME) coupling, between these two

properties is very attractive to researchers. Through ME coupling, dielectric

polarization variation can respond to an applied magnetic field, or magnetization can

be manipulated by an external electric field. Generally, several ME coupling

mechanism, such as strain/stress, interfacial charge, exchange coupling, can exist in

multiferroics material which contains magnetic and electrical phases.

Compared to single phase multiferroics, multiferroic composites, consisting of

separate ferroic phases with various connection schemes, usually display large

magnetoelectric coupling through magnetostrictive and piezoelectric effects.

Figure 1.1 The relationship between multiferroic and magnetoelectric materials - The

relationship between multiferroic and magnetoelectric materials

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1.2 Strain/Stress mediated ME coupling

Strain mediated magnetoelectric (ME) coupling in layered

ferromagnetic/ferroelectric heterostructures provides great opportunities in realizing

novel multiferroic devices, such as magnetoelectric random access memories

(MERAMs). Hu and coworkers [10] simulated the phase field and then demonstrated

a novel approach to voltage-controlled magnetic random access memory (MRAM).

They used the strain-mediated magnetoelectric coupling to control the direction of

magnetization in magnetic tunneling junction (MTJ) on a ferroelectric layer

heterostructure. A 90o rotation of the in-plane magnetization of the free layer can be

manipulated by strain mediated ME in the MTJ. This model of these voltage

controlled MRAM devices shows the ultra-low writing energy (less than 0.16 fJ per

bit), room temperature operation, high storage density, good thermal stability and fast

writing speed. Also, the voltage control of other magnetic properties:

magnetoresistence, exchange bias and magnetic domain wall propagation were also

studied experimentally by researches.

The modification of the magnetism by ferromagnetic phase shows a typical

“butterfly” like behavior as function of bipolar E-field in strain induced ME coupling.

This “butterfly” curve is due to the piezoelectricity of ferroelectric phase from

ferroelectric domain wall switching. However, the piezostrain at zero E-field is zero

resulting in volatile magnetization state. This will limit information storage or

MERAM devices, in which the magnetic state should be further controlled by voltage

impulses.

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Researchers demonstrated the non-volatile switching of magnetism in

ferromagnetic materials on different ferroelectric slab, such as (001) and (011)

oriented PMN-PT single crystal, (011) oriented PZN-PT single crystal and PZT

ceramic slab, experimentally. Wu et al. [11] realized a revisable and permanent

magnetic anisotropy reorientation in a muliferroic Ni/(011) oriented PMN-PT

heterostructure. They achieved a 300 Oe anisotropy field change in that system. The

change is non-volatile and is able to switch back and forth by E-field below coercive

field.

Figure 1.2 Normalized Kerr rotation hysteresis curves (M-H) along they y direction

under different electric fields (letters are the representatives of the labeled strain states

in the inset). The inset shows in-plane strain difference (εy−εx) as a function of

electric field. The drawings indicate the magnetization state: (c) permanent easy plane,

[(a) and (b)] temporary easy axis along x¯, and [(d) and (e)] permanent easy axis

along x¯.

As shown in Figure 1.2, the Kerr rotation hysteresis loop of Ni/PMN-PT in

application of different E-field leading to different piezo-strain state. The inset of

Figure 1.2 shows the relative strain difference as a function of E-field. By driving the

electric field from A to C (A-B-C), which represents the linear piezoelectric effect in

the linear ferroelectric regime. From the M-H hysteresis loop, the remnant

magnetization increases linearly in this area. However, in the linear regime, the

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magnetic anisotropy change is volatile and the magnetic state would go back to its

initial state after removing the E-field. When decreasing the E-field from 0 to coercive

field (D), the non-180° polarization reorientation dominates in PMN-PT and

introduces a sudden increase of strain into system. The large change of the magnetic

anisotropy was observed. When removing the E-field (E), the strain remains and the

magnetization is retained. As they increased the electric field from 0 to coercive field

(B), another non-180°polarization reorientation occurs back to the initial poling

direction. Thus the remnant strain is released, the magnetic state is switched back,

which confirmed the non-volatile switching. However, the mechanism behind that

using non-180° ferroelectric domain wall reorientation is unclear.

Most recently, Ming et al. [12] showed an unique ferroelastic switching pathway in

(011) oriented PMN-PT (0.71Pb(Mg1/3Nb2/3)O3-0.29PbTiO3) single crystal, which

allows up to 90% of polarization to rotate from an out-of-plane to a purely in-plane

direction (71o and 109

o polarization switching). They then produced two distinct,

stable and electrically reversible lattice strain states through this methos. Domain

distortion, polarization switching pathway and lattice strain responsing to in situ

perpendicular voltage in PMN-PT (011) are clearly presented using reciprocal space

mapping (RSM) and piezoforce microscopy (PFM) technology.

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Figure 1.3 (a) In-plane magnetic hysteresis loops of FeCoB/PMN-PT (011). Insets

are schematic (upper left) and FMR spectra (bottom right). (b) Schematic of FMR

measurement for (c-f). The sample is laid face down on an S-shape co-planar

waveguide. Magnetic fields are applied in the [100] direction and electric fields are

applied along the [011] direction. (c) Electric field dependence of the FMR frequency

in field sweeping mode. (d) Electric field dependence of the FMR field in frequency

sweeping mode. (e) FMR frequency responses under unipolar (red) and bipolar (blue)

sweeping of electric fields at room temperature. (f) Voltage-impulse-induced

non-volatile switching of FMR frequency.

See Figure 1.3, the multilayer films of Au(5 nm)/Fe60Co20B20(50 nm)/Ti(5 nm)

were deposited on (011) oriented single crystalline PMN-PT substrates and the

multiferroic heterostructure was characterized by E-field dependence of ferromagnetic

resonance field using coplanar-waveguide (CPW) FMR test system. Figure 1.3(e)

shows the resonance frequency dependence of E-fields. A "Butter-fly" curve (blue) is

observed as cycling triangle E-fields. With a positive E-field on a negatively poled

FeCoB/PMN-PT (011), a giant frequency jump happens near the coercive field around

1.5 kV cm-1

. When the polarization undergoes 71°and 109°ferroelastic switching

from the out-of-plane to the in-plane direction related to a lattice strain induced by the

domain distortion. Therefore, the hysteresis loop of the FMR frequency as a function

of the E-field is observed, Figure 1.3 (e). Like magnetic memory, two stable and

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reversible frequency remnant states A and B would facilitate the realization of

non-volatile frequency switching by reversing the E-field at the coercive field. The

voltage impulse induced magnetization switching was also realized shown in Figure

1.3 (f). As a PMN-PT (011) is subjected to an impulse of -6 kV cm-1

, the remnant

strain state A is retained and resulting in the largest FMR frequency of 9.9 GHz. Upon

applying an impulse field of 1.5 kV cm-1

, the resonance frequency is reduced to 7.6

GHz, indicating that the strain state is switched to B.

Figure 1.4 Schematics of domain structures and reciprocal space maps (RSMs) about

(022) and (002) reflections of PMN-PT (011) under various applied electric fields and

thus poling states. The first column (a,e,i) is for the unpoled state. The second column

(b,f,j) is for the positive poling state with up to 90% of polarization pointing upward.

The third column (c,g,k) is after applying an negative electric field of -1.5 kV cm-1

and then switching it off. The fourth column (d,h,l) is achieved by applying a positive

electric field of 5 kV cm-1

and then switching it off.

Futher, high resolution x-ray diffraction (HRXRD) measurements were used to

understand the polarization switching pathway and lattice strain in response to

E-fields. Figure 1.4 shows the E-field response to the reciprocal space maps (RSMs)

in the vicinity of the (022) and (002) reflections of the bare PMN-PT (011) substrates.

For the unpoled state of PMN-PT (011) (the first column in Figure1.4), a single broad

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spot is observed in both (022) and (002) reflections Figure 1.4(e) and (i). The analysis

of RSM patterns indicates that two possible domain structures r3 and r4 are dominant

in the unpoled state, and most of the polarization are in the plane. As the sample is

vertically poled with a strong positive voltage, the RSM in Figure 1.4(f) demonstrates

an addition high intensity (022) reflection spot with a lower Q022 value, corresponding

to the r1/r2 domain structures. At mean time, the intensity of the spot corresponding to

r3/r4 reduces dramatically. This represents that 71 ° and 109 ° ferroelastic

polarization switches from the in-plane direction to the out-of-plane direction

dominates and results in a large out-of-plane lattice strain. After applied a negative

E-field of -1.5 kV cm-1

is applied and removed, the domain distortion returns to r3/r4

and polarization is suppressed from the out-plane direction to the in-plane direction

Figure 1.4(g). As a large positive E-field of 5 kV cm-1

is applied and then switched off,

the domain structure is switched again and back to r1/r2, Figure 1.4(h). Thus, a stable

and reversible ferroelastic domain switching pathway is confirmed in their experiment,

which enables polarization rotation between the in-plane direction and the

out-of-plane direction.

1.3 Interfacial charge E-field tuning of magnetism in multiferroic

heterostructures

The strain/stress induced ME coupling suffered from substrate clamping effect,

which reduces its ME coupling strength, nevertheless, the charge mediated

magnetoelectric effect was first reported by Weisheit et al, does not limited by

substrate, especially in ultrathin magnetic layer. The magnetocrystalline anisotropy of

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ultra-thin iron-platinum and iron-palladium magnetic layer can be reversibly

controlled by E-field in an electrolyte in previous experiment, showing that the

screening charge provided by liquid electrolyte modified the intrinsic magnetic

properties. The one example of voltage control of magnetism offers an opportunity for

E-field induced magnetoresistance change in magnetic tunnel junctions (MTJ), the

core portion of MRAM devices. Maruyama et al. [13] also reported the change of

magnetic anisotropy in a Fe(001)/MgO(001) junction. With an E-field to dielectric

MgO layer, the surface magnetic anisotropies in 3d ferromagnetic metal/noble metal

interfaces were changed by the electron filling of 3d orbitals. From this mechanism,

they discovered a 40% change in the magnetic anisotropy by comparably small

E-field which could lead to varies application in low power spintronic devices.

In previous research, people also found the charge mediated ME coupling strength

is highly related to magnetic film thickness. For instance, ME coupling strength of

Fe/MgO heterostructure measured by Kerr hysteresis looper was significantly

dependent on Fe film thickness, at which, the maximum magnetic surface anisotropy

change was obtained at spin reorientation point. In Co20Fe80/MgO heterostructure,

magnetic surface anisotropy change decreased rapidly as Co20Fe80 film thicknesses

were larger than 0.5nm. Nevertheless, the mechanism causes charge mediated ME

coupling strength dependence on magnetic film thickness is still not certain. To

optimize the charge mediated ME coupling tunability in real applications, recently,

Zhou and Nan et al. [14] studied the voltage dependent ferromagnetic resonance

(FMR) in Ni0.81Fe0.19 (NiFe)/SrTiO3 (STO) magnetic/dielectric thin film

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heterostructures to quantitatively determine the thickness dependence of charge

mediated magnetoelectric coupling. Voltage induced FMR field change was carried

out through charge effect induced magnetic surface anisotropy change. Large voltage

induced FMR field shift of 65 Oe and magnetic surface anisotropy change of 5.6

kJ/m3 were obtained in NiFe/STO heterostructures. The voltage induced magnetic

surface anisotropy showed a strong dependence on the thickness of the magnetic thin

films, which was discussed based on the thin film growth model at the low thickness

side, and on the charge screening effect at large thickness side. The

thickness-dependent surface charge-mediated ME coupling has been studied in

bi-layered NiFe/STO thin film heterostructures with varied thicknesses of the NiFe

layer from 0.7 to 1.5 nm. High ME coupling induced FMR field shift of 65 Oe was

obtained and measured by ESR system, corresponding to large voltage tunable

effective magnetic anisotropy of 5.6 kJ/m3 and surface anisotropy of 6.7 μJ/m

2. This

investigation established a significant progress for magnetic/dielectric

heterostructure’s application in novel interfacial charge mediated magnetoelectric

devices. The detail discussion will be shown in Chapter 4.

For the real application by charge effect in magnetic tunnel junction, Wang et al.

[15] have demonstrated an electric-field-assisted switching in MgO based MTJ.

Compared with a traditional MTJ, a relatively thick ferromagnetic layer larger than 4

nm is used with in-plan or perpendicular magnetic anisotropy. But for enabling a

screening charge induced magnetoelectric effect, a ultra-thin ferromagnetic layer with

well defined metal/oxide interface is needed in the MTJ. In that work, the MTJ

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showed a TMR ratio of 118% with a core structure of

CoFeB(1.3nm)/MgO(1.4nm)/CoFeB(1.6nm), see Figure 1.5. They demonstrated that

the voltage controlled magnetocrystalline anisotropy can be used to switch the

magnetization of CoFeB from perpendicular to in-plane direction. The electric field is

used to reduce the HC of ferromagnets in MTJ and therefore the current that used to

switch the magnetization is two orders of magnitude less than the conventional one.

Figure 1.5 Electric-field-induced unipolar switching. (a) Normalized minorloops of

the TMR curve at different Vbias values. Inset: The full TMR curve at near-zero Vbias

where both ferromagnetic layers are switched by magnetic field. This MTJ has the

structure of CoFeB(1.3 nm)/MgO(1.2 nm)/CoFeB(1.6 nm). (b) Unipolar switching of

the MTJ by a series of negative pulses schematically shown in purple at the bottom)

with alternating amplitudes of -0.9 V and -1.5 V. The corresponding electric fields are

-0.75 V/nm and -1.25 V/nm, respectively. A constant biasing magnetic field of 55 Oe

in favour of the antiparallel state at -0.9 V was applied. (c) chematic diagram of the

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hysteresis loops of the top CoFeB layer showing the unipolar switching process:

magnetization-down - up switching at V = V1 (red) through STT with greatly reduced

energy barrier; magnetization-up ! down switching at V = V2 (black) by another

negative electric field, where│V2 >│V1│. The loop for V = 0 is shown in blue.

The vertical dotted line represents the position of the constant Hbias. The moment of

the bottom CoFeB is fixed pointing down.

Shoita et al. [16] also reported a coherent magnetization switching in a few atomic

layers of FeCo using voltage pulses, see Figure 1.6. They showed coherent

magnetization switching in ultra-thin MTJ by short voltage pulses of certain time

duration. FeCo layer was tilted from its initial in-plane magnetization to nearly

perpendicular orientation by bias magnetic field. The perpendicular anisotropy was

enhanced by E-field pulses produced a corresponding rotation of the magnetization

around the bias field. By applying the short-time E-field pulses, the magnetization

could be stopped at its original or the 90ᵒ with respect to this direction.

Figure 1.6 Macro-spin model simulation of coherent magnetization switching under

various pulse duration conditions. (a) Shape of the applied voltage pulse used in the

simulation. Pulse durations, pulse, are full-widths at half-maximum with rise and fall

times of 70 ps. (b) Examples of calculated trajectories induced by voltage pulse

application. Initial state (I.S.) and final state (F.S.) represent the magnetization state

before and after pulse voltage application.

1.4 Voltage control of magnetism in magnetic/BiFeO3 heterostructure

The room temperature single phase multiferroic, BiFeO3 (BFO), has attracted a lot

of recent research interest due to the coexistence of robust ferroelectricity (P) and

23

antiferromagnetism (L), and a weak canted magnetic moment (MC). In bulk BFO, the

weak moment results from the canting of the magnetic sublattices due to the

Dzyaloshinskii - Moriya (DM) interaction [17] as predicted by density functional

theory and confirmed experimentally [18]. E-field control of magnetism, like

magnetoresistance [19], magnetic anisotropy and magnetization, in a ferromagnetic

layer exchange coupled to BFO layer has been most recently reported. Heron et al.

[19] discovered a nonvolatile, room temperature magnetization reversal determined

by an electric field in a CoFe/BFO multiferroic heterostructure. Figure 1.7 revealed

that there is an one-to-one corelation between stripe-like ferromagnetic domain in

CoFe and ferroelectric domain in BFO, resulting in an uniaxial magnetic anisotropy

of CoFe, Figure 1.7(a). After applied a voltage across BFO layer, a magnetization

reversal was confirmed by anisotropic magnetoresistance (AMR) measurements. This

experiement give the evidence of the coupling between CoFe magnetic moment and

canted moment in BFO, see Figure 1.8. More detail research regarding to how the

canted moment switching with electric field was discussed in section 4.

24

Figure 1.7 (color online). (a) In-plane M-H curves measured every 45� at room

temperature from CoFe/BFO heterostructures. The CoFe growth field was applied

along (black open circles) or perpendicular [gray (red) open circles] to the net inplane

polarization direction (Pnet IP). (b) In-plane PFM image of BFO. (c) XMCD-PEEM

image of the CoFe=BFO heterostructure. The gray (blue) and black arrows in (b) and

(c) correspond to the in-plane projections of the polarizations in each of the

ferroelectric domains of BFO and to the magnetic moments in the CoFe layer,

respectively.

25

Figure 1.8 (color online). (a) Open black circles show the high field (2000 Oe) AMR

response (top panel). The low-field (20 Oe) AMR response for the as-grown state is

plotted with the open red circles (second panel from top). The open blue circles show

the low-field AMR after pulsing an electric field of 130 kV cm-1

in zero magnetic

field (second panel from bottom). Application of a -130 kV cm-1

electric-field pulse

results in the recovery of the phase of the as-grown low-field AMR response (open

green circles, bottom panel). (b),(c) Representations of the one-to-one magnetic

interface coupling in the CoFe/BFO heterostructure in the (b) as-grown state and (c)

after the first electric pulse.

26

Chapter 2 Magnetic/Ferroelectric Multiferroics for Tunable

Microwave Applications

2.1 Low moment approach of ME coupling in NiCr/ferroelectric

multiferroics heterostructure

2.1.1 NiCr/PZT and NiCr/PZNPT multiferroic heterostructure fabrication

Layered magnetic/piezoelectric multiferroic heterostructures such as

FeGaB/PZN-PT, [20] Fe3O4/PZN-PT, [21] with a magnetic thin film on piezoelectric

slab provides a great opportunity to achieve strong ME coupling. The E-field induced

effective magnetic field of magnetic film on a ferroelectric slab can be described by

the formula of: SeffSeff MYEdH /3 , [20] in which, λs is the magnetostriction, Y

represents Young's modulus, Ms is the saturation magnetization of the magnetic thin

film, deff and E are the effective piezoelectric coefficient and applied E-field on the

ferroelectric slab, respectively, and v is Poisson ratio of the NiCr film. Increasing the

magnetostriction and/or reducing the saturation magnetization would be two

approaches to achieve strong E-field induced magnetic field Heff. Lou et al reported

new RF FeGaB films with a large magnetostriction coefficient of 70 ppm, a giant

piezomagnetic coefficient of 7 ppm/Oe, and a saturation magnetization of 1.4 Tesla,14

and demonstrated a large electric field induced magnetic field of 750 Oe in

FeGaB/PZN-PT (lead zinc niobate lead titanate) heterostructures [20].

The NiCr alloy system has low saturation magnetization and relatively high

magnetostriction, which can be a good candidate for low moment multiferroic

27

heterostructures. In this paper, we investigated NiCr alloy thin films with different Cr

contents, which showed a low magnetization of 1100~1910 Gauss and a relatively

high magnetostriction of -5.1 ~ -7.8 ppm. The low magnetization and high

magnetostriction in NiCr alloy films lead to a high ME coupling coefficient of 13

Oe·cm/kV (NiCr/PZT) and 75.6 Oe·cm/kV (NiCr/PZN-PT)), compared with

FeGaB/Si/PZT (2 Oe·cm/kV), [22] FeGaB/PZN-PT (86 Oe·cm/kV), [20]

Fe3O4/PZNPT (108 Oe·cm/kV), [23] Zn0.1Fe2.9O4/PZN-PT (23 Oe·cm/kV). [24]

The frequency tunability is 39 MHz·cm/kV (NiCr/PZT) and 250 MHz·cm/kV

(NiCr/PZN-PT). Hence, these the NiCr/PZT and NiCr/PZN-PT heterostructures with

strong magnetoelectric coupling have great technological potential.NiCr alloy

magnetic thin films were deposited by the DC magnetron co-sputtering with Ni and

Cr targets at room temperature on Si substrates with different Ni/Cr ratios. All films

were deposited for 600 seconds, leading to a film thickness of ~50 nm. NiCr

compositions were measured by X-ray fluorescence (XRF) system. The static NiCr

hysteresis loops with different components were measured by vibrating sample

magnetometer (VSM). Microwave ME interaction was investigated by a broadband

ferromagnetic resonance spectrometer. Static electric field was applied across the

NiCr/PZT and NiCr/PZN-PT samples thickness direction for achieving electric field

tuning of the magnetic properties.

2.1.2 Magnetic Properties of NiCr thin films

Figure 2.1 (a) shows the out-of-plane hysteresis loops of Ni1-xCrx with different Cr

contents x. We can observe a clear trend that the saturation magnetization gradually

28

decreases as Cr content increases. At x=0.046, 0.05, 0.054, 0.059, 0.061 the

out-of-plane hysteresis loops exhibit characters of a ferromagnetic material with a

non-zero remnant magnetization with a saturation magnetization of 1910 Gauss, 1550

Gauss, and 1100 Gauss, 1030 Gauss, 820 Gauss, respectively. As the x increased to

0.064 or higher, the NiCr magnetic thin film starts to show signs of being

superparamagnetic at room temperature with zero remnant magnetization and zero

coercivity.

Figure 2.1 (a) Out-plane Hysteresis loop of Ni1-xCrx alloy thin film on Si substrate

with different Cr content x on left hand side. (b) In-plane hysteresis loop of Ni1-xCrx

alloy thin film on Si substrate with different Cr content x on right hand side. Both

M(H) loops are measured at room temperature.

Figure 2.1 (b) shows the in-plane hysteresis loops of Ni1-xCrx thin films with

different Cr contents x. The hysteresis loops of x=0.046, 0.05, 0.054 exhibits typical

out-of-plane magnetization, implying that there may exist an magnetoelastic

anisotropy (E = (3/2)ζλ) associated with a tensile stress and a negative

magnetostriction of the NiCr film. Figure 2.2 shows the magnetic field sweep FMR

spectra of the NiCr alloy films measured at 11.3 GHz with field sweep range from

29

500 Oe to 3000 Oe. We can only obtain clear a FMR signal of Ni1-xCrx thin films

with x=0.054, 0.05, 0.046 but we could not see the FMR spectrums in Ni1-xCrx

alloys with x=0.59 and larger. We did not show their magnetic properties in table I.

The resonance field was 2740 Oe at x=0.054, 2490 Oe at x=0.05 and 2250 Oe at

x=0.046. The FMR linewidth was about 250 Oe for all three films. From the

Landau–Lifshitz equation, [20] )4)(( sresaresares MHHHHf , where ɣ is

gyromagnetic constant of 2.8 MHz/Oe, we can calculate Ha from measured 4πMS, fres

and Hres.

Figure 2.2. In-plane field-sweep ferromagnetic resonance spectra of Ni1-xCrx films

with different Cr content x measured at 11.3 GHz

Saturation magnetostriction values of thin films with varied Cr contents were

estimated by the electric field induced effective magnetic field through FMR field

shift. The FMR field shifts of the Ni1-xCrx/PZN-PT heterostructures are 260 Oe, 228

Oe and 211 Oe, corresponding to Cr contents of x=0.054, 0.05, 0.046, respectively.

By substituting parameter, d31 (-3000 pC/N), d32 (1100 pC/N), Young's Modulus of Ni

thin film (93 GPa) into ME coupling formula: SeffSeff MYEdH /3 , where

deff=(d31-d32)/(1+v), [20] the magnetostriction of Ni1-xCrx films can be calculated as

30

-5.1 ppm, -6.3 ppm and -7.2 ppm with a Cr content of x=0.046, 0.05, 0.054. The

Ni1-xCrx thin film with a Cr content x=0.054 (All NiCr films listed below represent

Ni0.946Cr0.054 films) and largest ME coupling coefficient, which was chosen for further

investigation on both PZT (d31=-400 pC/N) and PZN-PT substrates.

It is important to choose a high ratio of saturation magnetostriction over saturation

magnetization (λs/Ms) based on the formula: SeffSeff MYEdH /3 , to achieve high

magnetoelectric coupling coefficient. Ni0.946Cr0.054 film is a good choice for an

investigation into ME couplingwhich shows a large λs/Ms ratio with a saturation

magnetostriction of -5.1 ppm and saturation magnetization of 1100 Gauss. NiCr

(Ni0.946Cr0.054) thin films were deposited on polished PZT substrates and (011) cut

PZN-PT single crystal slabs. The dimensions of these substrate are 1 cm×0.2 cm×0.5

mm (PZT) and 1 cm (100)×0.5 cm (01-1)×0.5mm (PZN-PT). The NiCr top layer had

a thickness of 85 nm and the thickness of Cr electrode bottom layer was 100 nm. In

this experiment, we applied a high voltage from 400 V to -600 V on the NiCr/PZT

heterostructure, which corresponds to an electric field of 8 kV/cm to -12 kV/cm and

also a high voltage from -100 V to 400 V on NiCr/PZN-PT multiferroic

heterostructure with electric field tunable range of ~ -2 ~ 8 kV/cm.

31

Figure 2.3 (a) Electric field dependence of the in-plane magnetic field sweep FMR

spectra of NiCr/PZT multiferroic heterostructures measured at 6.85 GHz. (b) Butterfly

plot of anisotropy magnetic field as a function of applied electric field from -8 kV/cm

to 12 kV/cm. (c) Electric field dependence of the in-plane frequency sweep FMR

spectra of NiCr/PZT multiferroic heterostructure measured at 50 Oe. (d) Butterfly plot

of resonance frequency as a function of applied electric field form -8 kV/cm to 12

kV/cm.

2.1.3 Electric field control of magnetism in NiCr/PZT and NiCr/PZNPT

heterostructure

The electric field controllable FMR behavior of the NiCr/PZT multiferroic

heterostructures at a given resonance frequency of 6.85 GHz was measured on an

FMR spectrometer, and is shown in Figure 2.3 (a)~(d). The field sweep FMR spectra

in Figure 2.3(a) exhibited E-field controllable resonance magnetic field under

different applied E-field from -12 kV/cm to 8 kV/cm. The external magnetic field is

applied parallel to the long axis (1 cm) direction of the PZT substrate and the E-field

32

was applied along PZT thickness direction, from NiCr thin film top layer to Cr

electrode bottom layer. There was a high resonance magnetic field shift from 1034 Oe

to 1294 Oe, or an effective magnetic field of 260 Oe. Figure 3 (b) demonstrates the

butterfly behavior of E-field control of effective anisotropy fields with E-field varied

from -12 kV/cm to 12 kV/cm. By fixing the magnetic bias field at 50 Oe, the FMR

measurement system can be also used to measure frequency sweep spectra for

NiCr/PZT multiferroic structure as shown in Figure 2.3 (c). It is clear that a large

FMR shift of 0.78 GHz (from 4.062 GHz to 3.282 GHz), or fmax/fmin= 1.24 was

achieved by applying E-field varied from -12 kV/cm to 8 kV/cm. Figure 2.3 (d)

represents the butterfly curve of resonance frequency and E-field (-12 kV/cm to 12

kV/cm) of the NiCr/PZT heterostructure, which exhibits a linear dependence between

tunable FMR frequency and the electric field when E is less than Ecritical.

33

Figure 2.4. (a) Electric field dependence of the in-plane magnetic field sweep FMR

spectra of NiCr/PZN-PT multiferroic heterostructures measured at 6.85 GHz. (b)

Butterfly plot of anisotropy magnetic field as a function of applied electric field form

-2 kV/cm to 8 kV/cm. (c) Electric field dependence of the in-plane frequency sweep

FMR spectra of NiCr/PZN-PT multiferroic heterostructure measured at 50 Oe. (d)

Butterfly plot of resonance frequency as a function of applied electric field form -2

kV/cm to 8 kV/cm.

Figure 2.4 (a) shows the field-sweep FMR behavior of the NiCr/PZN-PT

multiferroic heterostructure at a given resonance frequency of 6.85 GHz on the FMR

system, similar to Figure 2.4(a) for the NiCr/PZT multiferroic heterostructure. The

external magnetic field is applied parallel to the in-plane [011] direction of the

PZN-PT single crystal and the E-field is applied along PZN-PT thickness direction,

from the NiCr thin film top layer to the Cr electrode bottom layer. The resonance

34

magnetic field was shifted from 1171 Oe to 1927 Oe under different applied E-fields

from -2 kV/cm to 8 kV/cm, corresponding to a giant magnetic resonance field shift of

756 Oe and a large magnetoelectric coupling coefficient of dH/dE= 75.6 Oe cm/kV.

Figure 2.4 (b) shows the butterfly behavior of E-field tunable anisotropy field, and

Figure 2.4 (c) shows frequency sweep spectra for the NiCr/PZT multiferroic structure

under an applied bias magnetic bias field of 50 Oe. A large resonance frequency shift

from 1.271 GHz to 3.771 GHz was achieved by changing the E-field from -2 kV/cm

to 8 kV/cm, corresponding to fmax/fmin =2.97, or 250 MHz cm/kV of tunable frequency

range. Figure 2.4 (d) demonstrates the butterfly behavior curve of resonance

frequency and E-field (-8 kV/cm to 8 kV/cm).

Figure 2.5. (a) E-field dependence of NiCr Gilbert damping constant on PZT and

PZN-PT substrates. (b) E-field dependence of NiCr ∆H0 on PZT and PZN-PT

substrates.

The Gilbert damping coefficients of the NiCr films in NiCr/PZT and NiCr/PZN-PT

heterostructures were measured at different bias E fields, which were extracted by

using the following equation α=0.5ɣ·(∆H-∆H0)/f0, where ɣ is the gyromagnetic

constant ~2.8MHz/Oe, ∆H is the FMR linewidth, ∆H0 is the intercept of y-axis

35

linewidth and f0 is the FMR frequency. We measured 6~7 Oe NiCr FMR ∆H

linewidths under different FMR frequencies, f0, at certain E-field and then

calculated the linear equation, ∆H=2αf0/ɣ+∆H0, between ∆H and f0 by doing linear

extrapolation. The α and ∆H0 at that E-field can be obtained through the slope and the

y-axis intercept of the linear equation, correspondingly. Figure 2.5 (a) shows the

Gilbert damping coefficients of NiCr thin films on PZT and PZN-PT substrates as a

function of the E-field applied on piezoelectric substrates. The Gilbert damping

coefficients increase monotonically from 0.0072 at -12 kV/cm to 0.0078 at 8 kV/cm

for NiCr/PZT; and from 0.0072 at -2 kV/cm to 0.0086 at 8 kV/cm for NiCr/PZN-PT.

Figure 2.5 (b) demonstrates the ∆H0 dependence of E-fieldwhich varied similar to the

Gilbert damping constant. ∆H0 increase from 190 Oe at -12 kV/cm to 203Oe at 8

kV/cm for NiCr/PZT; and from 193 Oe at -2 kV/cm to 216 Oe at 8 kV/cm for

NiCr/PZN-PT. The E-field dependence of the Gilbert damping coefficients and ∆H0

can be explained by the E-field induced effective magnetic field, which constitutes

added benefits for E-field tunable RF/microwave magnetic devices.

It is worth noting the ME coupling coefficient of NiCr/PZN-PT is slightly lower

than our previous results demonstrated in Fe3O4/PZN-PT and FeGaB/PZN-PT. That is

because the λs of NiCr is decreased as 4πMs decreases with Cr doping, which leads to

a relative smaller λs/Ms ratio. Higher λs/Ms ratio of 1.7 ppm/Gauss was discovered in

Terfenol based alloys, such as Tb1-xNdx(Fe0.9B0.1)2 alloys. [25] However, these

terfenol based alloys typically have very large FMR linewidth and are also very

expensive. This investigation on NiCr films and the multiferroic heterostructures

36

based on NiCr alloys constitutes the first attempt to develop magnetic materials with

low moments and high magnetostrictions in order to achieve higher ME coupling

coefficient.Future efforts on low-moment magnetic films for multiferroics should put

more emphasis on achieving a high λs/Ms ratio while maintaining good RF properties

at a reasonable cost.

2.2 Non volatile tunable FeGaB/PSZT magnetic/antiferroelectric

heterostructures

2.2.1 FeGaB/PSZT multiferroic heterostructure fabrication

Besides tunability, there are still challenges exist in these ME devices. For example,

tunability and volatility are critical properties in voltage-tunable RF/microwave ME

devices, such as tunable filters and resonators [1-6]. Currently, many ME devices

require a constant applied E-field rather than a short time voltage impulse for tuning

and manipulation. Driving by the motivation of reducing energy consumption, the

non-volatile voltage impulse tunable ME devices, while at the same time enabling

large and distinct E-field manipulating magnetic properties, such as, magnetization,

ferromagnetic resonance(FMR), etc, was investigated.

In this work, we reported novel magnetic/antiferroelectric heterostructures of

amorphous FeGaB film on La-modified Pb(Sn,Zr,Ti)O3 (PSZT) ceramic substrates.

The FeGaB films were deposited on the top or on the side of the antiferroelectric

PSZT substrate by physical vapor deposition (PVD) system (Figure 2.6(a)). We

37

systematically studied magnetic/microwave performance in FeGaB/PSZT

multiferroics heterostructure under varying E-field. Strong ME coupling of ~80 Oe

was exhibiting in FMR field measurements, which can be generated in engineering

requirements. Mostly importantly, by introducing E-field induced

anti-ferroelectric/ferroelectric phase transition of PSZT into multiferroics system, a

novel non-volatile tuning magnetic/microwave properties induced by voltage impulse

can be achieved in FeGaB/PSZT system. The strong magnetoelectric coupling with

voltage impulse tunable non-volatile switch in FeGaB/PSZT

magnetoelectric/antiferroelectric heterostructures constitutes a novel approach to

achieving strong magnetoelectric coupling which can have great technological

implications.

Multiferroic heterostructure FeGaB/PSZT are prepared by co-sputtering of Fe70Ga30

and B targets onto La-modified PSZT ceramic substrates (8 mm Length×3 mm

Width×0.5 mm Height ) with a base pressure below 1×10−7

Torr at room temperature.

The La-modified PSZT ceramics Pb0.96La0.04(Zr0.45Sn0.36Ti0.18)O3 substrates were

prepared by a conventional solid-state reaction process. Raw powders were mixed

with Al2O3 balls in deionized water by ball-milling for 2 hours. The mixtures were

calcined at 850 oC for 2 hours after being dried. After ball-milling, the powder was

pressed into disks. Finally, the green compacts were sintered at 1340 oC for 2 hours in

lead ambiance. The surfaces were polished to deposit magnetic 100nm FeGaB thin

film and 50 nm Cr electrodes. A 5-nm-thick Cr layer was inserted between FeGaB

layer and PSZT ceramics to improve adhesion. The ferroelectric/antiferroelectric and

38

piezo-strain properties were measured by P(E) loop and Photonic meter(MTI 2000).

Ferromagnetic resonance field sweeping measurements were carried out by electron

spin resonance(ESR) measurement. A DC E-field was applied across the thickness

direction of PSZT coated with Cr electrode on the back as an electrode. The

magnetization measurements of FeGaB/PSZT were carried out by using a vibrating

sample magnetometer(VSM) (Lakeshore 7400).

Figure 2.6 (a) The schematic of FeGaB film deposited on the top or on the side of

Pb(Sn,Zr,Ti)O3 ceramics. E-field is applied across PSZT layer; (b) X-ray diffraction

pattern of PSZT ceramics (c) Polarization and strain vs. E-field loop of

Pb(Sn,Zr,Ti)O3 ceramic material, correspondingly; (d) Strain dependence of E-field,

from 0 kV/cm to 30 kV/cm.

The X-ray diffraction(XRD) pattern of PSZT ceramics was measured with a Cu Ka

source (λ=1.541Å), see Figure 2.6(b), the typical PZT crystal orientations were

39

obtained in XRD measurements. The polarization vs applied E-field(P-E) loop shows

a typical antiferroelectric P-E loop, [26-29] which indicates the anti-ferroelectric

phase of PSZT ceramic, see Figure 2.6(c). As the applied E-field is larger than 20

kV/cm, associated with the polarization increases from from 0 to 18 μC/cm2, the

anti-ferroelectric phase is transferring into ferroelectric phase, leading to a large

E-field induced strain along d33 the side orientation, see Figure 2.6 (a) [26-31] as

shown in Figure 2.6(c). The typical structure of antiferroelectric lead zirconate

Pb(Zr,Ti)O3 system is orthorhombic (pseudo-tetragonal) [30] and it can be

voltage-induced into a rhombohedral ferroelectric phase [31]. A large strain was

approached because the c-axis is elongated during the transition. The

antiferroelectric-ferroelectric phase transition in PZST substrates gives us the

opportunity to obtain a strong magnetoelectric coupling coefficient due to large strain

change at phase transition point, furthermore, the hysteretic strain dependence of

E-field21-24 also provides the possibility of voltage impulse induced non-volatile

switch [32]. As demonstrated in Figure 2.6(d), E-field increased from 0 kV/cm to 30

kV/cm and then decreased to 0 kV/cm, the strain dependence of E-field follow an

identical hysteretic behavior induced by ferroelectric/antiferroelectric phase transition

of PSZT substrates. There is a large strain gap between the 15 kV/cm(green) E-field

increased from 0 kV/cm and the 15 kV/cm(blue) E-field decreased from 30 kV/cm,

which offers the non-volatility and controllability induced by voltage impulse.

40

2.2.2 Non-volatile control of magnetism in FeGaB/PSZT multiferroic

heterostructure

Figure 2.7 (a) (b) shows the magnetization vs applied magnetic field

hysteresis(M-H) loops measured under varying E-field, 100 nm FeGaB thin film was

prepared on the top (a) or the side (b) of PSZT substrates. The strain was larger on the

side of PSZT than the top of PSZT, however, the roughness was also large on the side

of PSZT than the top of PSZT. We studied both cases in our experiment to obtain the

optimized tunability and non-volatility in controlling the magnetization or FMR field.

In Figure 2.7(a) (b), the M-H loops dependence of applied E-field is studied. The

coercivity HC of the FeGaB film on the top surface of the PSZT substrate was

increased from 35 Oe to 41 Oe by applying an electric field, as represent Figure 2.7(a).

On the contrary, the HC of the FeGaB film on the side was decreased from 39 Oe to

27 Oe while an electric field of 30 kV/cm is applied, see Figure 2.7(b). At the same

time, the remanent magnetization of the FeGaB film on the side of PSZT was reduced

by 30% at an applied E field of 30 kV/cm at zero magnetic field. Further, we

examined the non-volatility from M-H loops, for FeGaB(top)/PSZT heterostructure,

the FeGaB M-H loop(green) measured at 15 kV/cm E-field increased from 0 kV/cm is

closed to the M-H loop(black) measured at 0 kV/cm E-field. Similarly, the M-H

loop(blue) measured at 15 kV/cm E-field decreased from 30 kV/cm is closed to the

M-H loop(red) measured at 30 kV/cm E-field. There exists a significant gap between

the two applied E-field of 15 kV/cm back and forth, introducing non-volatile

magnetization switches. At applied magnetic field(H=40.5 Oe), by switching the

41

E-field, the magnetization was changed from 500 Gauss to -500 Gauss, the

ΔM/M=17%, see the upper left inset of Figure 2.7(a). For FeGaB(side)/PSZT

heterostructure, as represented in Figure 2.7(b), non-volatile E-field induced M-H

loops switching was also obtained. The largest magnetization switches back and forth

were achieved at remnant magnetization, where ΔM=50 Gauss.

Figure 2.7 (a) M-H loops under varying E-field of FeGaB(top)/PSZT multiferroics

heterostructure; (b) M-H loops of FeGaB(side)/PSZT multiferroics heterostructure; (c)

FMR spectra under varying E-field of FeGaB(top)/PSZT multiferroics heterostructure;

(d) MR spectra of FeGaB(side)/PSZT multiferroics heterostructure.

Figure 2.7(c) (d) demonstrated ferromagnetic resonance field spectrums of top and

side FeGaB/PSZT respectively, under varying E-field. In FeGaB(top)/PSZT

heterostructure, the maximum FMR field switch is 32 Oe, from E-field of 0kV/cm to

42

30kV/cm, corresponding to ME coupling coefficient 1.1 Oe cm/kV. And the

maximum FMR field switch was 81 Oe, from E-field of 0 kV/cm to 30 kV/cm,

leading to ME coupling coefficient of 2.7 Oe cm/kV, at FeGaB(side)/PSZT

heterostructure. The electric field induced in-plane anisotropy field change can be

simulated by piezoelectric and inverse magneto-elastic equations. In our case, the

thicknesses of the FeGaB film and electrode layers are much less than that of the

PSZT single-crystal substrate, the FeGaB film experienced an in-plane stress induced

by the piezoelectric strain of the PSZT ceramics. As shown in the inset of Figure 2.7(c)

(d), the FMR field Hr dependence of applied E-field is similar to the strain

dependence of E-field, see Figure 2.6(b) (c), which can be derived from equation

ΔHr=ΔHeff=3λsεY/MS [9-10]. The FMR field shift is directly proportional to the

E-field induced strain, here Y is Young’s modulus of the FeGaB film, ε is the effective

piezo-strain along d33 direction, see Figure 2.6(b) (c), Y is Young’s modulus of FeGaB

and Ms is the saturation magnetization. For our FeGaB thin film, Y=55 GPa, λs=70

ppm, Ms=1.3 Tesla and ε is 0.07% for PSZT slab at applied E-field of 30 kV/cm, see

Figure 2.6(b), (c). Effective magnetic field ΔHeff can be calculated as 84 Oe, which

confirmed our experimental result of 81 Oe. The reason why FMR field shift of top

FeGaB/PSZT is much smaller than that of side FeGaB is d31 is smaller than d33(about

50%) of PSZT. The theoretical result of top FeGaB FMR field shift is 42 Oe, which is

close to 31 Oe as we measured in experiment.

43

Figure 2.8 (a) Magnetization switch of FeGaB(top)/PSZT under applied H-field of

40.5 Oe induced by voltage impulse; (b) Magnetization switch of FeGaB(side)/PSZT

under zero bias magnetic field induced by voltage impulse; (c) FMR field switch of

FeGaB(top)/PSZT induced by voltage impulse; (d) FMR field switch of

FeGaB(side)/PSZT induced by voltage impulse.

Based on the E-field induced non-volatile switches of magnetization at bias applied

magnetic field and FMR field in FeGaB/PSZT heterostructure, as demonstrated on

Figure 2.8, the voltage impulse(100ms) tunable magnetization and FMR field

mechanism can be designed. Figure 2.8(a) (b) shows the voltage impulse tuned

FeGaB magnetization at bias magnetic field, by maintaining a constant E-field of

15kV/cm, the E-field impulse(

44

E-field, see Figure 2.7(a) inset. Figure 2.8(b) showed the magnetization switches at

zero bias magnetic field in FeGaB(side)/PSZT induced by voltage impulse, from 113

Gauss to 50 Gauss, back and forth, see Figure 2(b) inset. The FMR field switch by

voltage impulse was also investigated. For FeGaB(top)/PSZT heterostructure, the

FMR field switches from ~1015 Oe to ~995 Oe by applying voltage impulse, as

demonstrated in Figure 2.8(c) and the FMR field switch from ~2094 Oe to ~2043 Oe

under same voltage impulse series, see Figure 2.8(d). The result also accorded with

the FMR field dependence of E-field measurements, see Figure 2.7(c) (d).

In summary, we have demonstrated large magnetic/microwave tunability through

E-field strain-induced ME coupling in FeGaB/PSZT multiferroics composites. A

non-volatile magnetization and FMR field switching by E-field-induced

antiferroelectric/ferroelectric phase transition in PSZT was realized. These features,

including large tunability and non-volatile switching gives FeGaB/PSZT

heterostructure great candidates for next-generation voltage-impulse-controlled

lightweight, energy efficient, spintronics RF/microwave devices.

2.3 Non volatile tunable FeGaB/PZNPT magnetic/ferroelectric

heterostructures with giant tunability

2.3.1 FeGaB/PZNPT multiferroic heterostructure characteration

In order to improve non-volatile control of magnetism with larger ME coupling

strength, recently, we reported preliminary results on a novel microwave

45

heterostructure of FeGaB/PZN-PT (Lead Zinc Niobate-Lead Titanate), 错误!未找到

引用源。 showing a large E-field-induced ferromagnetic resonance (FMR) tunable

range with a small line-width, which is ideal for microwave applications. In this work,

we systematically studied E-field control of microwave performance in the manner of

magnetic field sweeping and frequency sweeping in ME composites FeGaB/PZN-PT.

A strong ME interaction was demonstrated by a large E-field-induced in-plane strain

measured through in situ x-ray diffraction and verified by E-field tuning of FMR

field and frequency. A new technical solution with dual E-and H-field tunability was

developed to dramatically enhance FMR tunable range up to 13.1 GHz, which would

greatly satisfy the engineering requirements for different applications. In addition,

regarding the hysteric and irreversible E-field-induced phase transition in single

crystal PZN-PT substrate, we successfully realized novel voltage-impulse-induced

memory-type of magnetization switching and FMR tuning in FeGaB/PZN-PT

multiferroic heterostructures. An extremely large converse magnetoelectric coupling

coefficients were also demonstrated, which were 3850 Oe·cm kV-1

(∆H/∆E), 3620

Oe·cm kV-1

(∆H/∆E) at phase transition points of 3 kV cm-1

and 5.8 kV cm-1

,

respectively. The giant voltage tunable FMR frequency and voltage impulse induced

non-volatile magnetization switching in FeGaB/PZN-PT show great potential for next

generation RF devices with compact size, light weight and high energy efficiency.

46

Figure 2.9 (a) X-ray diffraction pattern of amorphous FeGaB on (011) orientated

PZN-PT substrate. Left inset is the AFM image of PZN-PT substrate, showing a

ferroelectric multi-domain. In each domain, the surface roughness is less than 0.5 nm.

(b) E-field induced lattice change in FeGaB/PZN-PT heterostructures along different

orientations. Insets at up-corner show lattices change along [100] and [111] directions

as electric field applied. Inset at right corner shows the diffraction pattern shift as

electric field applied. The overall displacement ratio is about -0.36% along [100] and

0.25% along [011].

(011) oriented single crystal PZN-PT with large in-plane piezoelectric coefficients

of d31=-3000 pC N-1

[100] and d32=1500 pC N-1

[01-1] was used as ferroelectric

47

substrate to obtain maximum electric-field-induced in-plane biaxial strain 错误!未找

到引用源。. Surface morphology of the PZN-PT substrate was characterized by

Atomic Force Microscope (AFM) in taping model as shown in left inset in Figure

2.9(a), exhibiting typical ferroelectric rhombohedral domains with kinks at domain

wall. Such domain structure is caused by the distortion from cubic to rhombohedral as

lowering temperature, which makes spontaneous polarization along the body

diagonals in pseduocubic PZN-PT cell. Within a single domain, the surface shows a

root mean square (rms) roughness of 0.55 nm. The film structures, as well as

voltage-induced strain or lattice changes in PZN-PT were characterized by in situ

x-ray diffraction (XRD). As shown in Figure 2.9(a), there are no peaks observed from

the film, except the peaks from the (011) oriented PZN-PT, indicating that amorphous

FeGaB phase was produced with excellent soft magnetism and narrow FMR

line-width. [33] The film thickness was determined to be 50 nm by fitting x-ray

reflectivity spectrum. To investigate electric-field-induced lattice change along

various orientations and estimate overall in-plane biaxial strain, we performed in situ

x-ray diffraction measurements on FeGaB/PZN-PT (011) structure under various

electric fields. Note that PZN-PT substrate is initially in poling state and electric field

was applied perpendicularly. An expansion along the out-of-plane direction associated

with an effective in-plane contraction was observed as the electric field was applied.

The enhanced out-of-plane lattice parameters of PZN-PT is visible as a shift of Δc/c =

+ 0.28 % as shown in figure 2.9(b) inset. Based on the lattice changes along various

orientations under electric fields, in-plane biaxial strain was calculated to be -0.36%

48

along [100] and 0.25% along [01-1] which are proportional to the in-plane

piezoelectric coefficients. As the FeGaB film is thin enough compared to the PZN-PT

slab, the FeGaB film on PZN-PT substrate will experience the same strain states as

the PZN-PT under an electric field, which has been proven in the following discussion

on electric field control of magnetic properties.

2.3.2 RF/microwave tunability of FeGaB/PZNPT heterostructure

Figure 2.10 E-field tuning FMR in field sweeping (a, b) and frequency sweeping (c,d)

showing a giant tunable range up to 1200 Oe and 5.3 GHz respectively.

49

Electric field tuning of microwave performance for both frequency sweeping and

field sweeping in FeGaB/PZN-PT were carried out in our homemade

coplanar-waveguide (CPW) FMR test unit. FeGaB/PZN-PT is laid face down on

CPW and magnetoelectrically operates in the L-T (Longitudinal

magnetized/Transverse polarized) mode, where voltage is applied along the normal

direction and in-plane magnetic anisotropy is manipulated by biaxial stress through

piezoelectric and magnetostrictive effects. As a function of magnetic anisotropy,

in-plane FMR of FeGaB can be expressed by Kittel equation

f =g (H +Heff )(H +Heff + 4pMs ) , where γ is the gyromagnetic ratio ~2.8 MHz Oe-1,

H is the FMR field; 4πMs is the magnetization of FeGaB. Heff is the voltage

induced effective magnetic field and can be expressed by Heff =

3lss EM s . Here, λ is

magnetostriction constant of FeGaB (~80 ppm); ζE is electric-field-induced biaxial

stress which is tensile along [01-1] and compressive along [100]. This biaxial stress,

which has been demonstrated in situ x-ray diffraction measurements, enables

electrically manipulating FMR in both frequency and field sweeping measurements.

As shown in Figure 2.10(a,b), remarkable upward and downward shifts in FMR field

spectra were observed in field sweeping as an electric field of 6kV cm-1 was applied,

while external magnetic fields are along [100] and [01-1] direction respectively,

confirming the result in ref. 34. Such opposite shift originates from the

electric-field-induced magnetic easy axis along [01-1] and hard axis along [100]. The

total FMR tunable range was 1200 Oe, corresponding to a large ME coefficient of 100

Oe cm kV-1

, indicating a strong mechanical coupling at interface between substrate

50

and thin film. Similar results were also observed in frequency sweeping FMR

measurement as shown in Fig 2.10(c,d). A downward frequency shift from 11.3

GHz to 8.5 GHz and upward shift from 11.3 GHz to 13.8 GHz achieved as an external

magnetic bias field of 518 Oe was applied along [100] and [01-1] direction

respectively, corresponding to a giant frequency tunable range of 5.3 GHz and

microwave ME coefficient of 880 MHz cm kV-1. These electric-field-induced FMR

field and frequency shifts are correlated and can be interpreted in Kittle equation (1).

In addition, FMR linewidth for both frequency sweeping and field sweeping shows

relative small and less than 60 Oe or 360 MHz no matter electric field applied,

indicating a uniform deformation occurred in FeGaB film under biaxial stress. We

also observed non-linear FMR response as electric field applied from 4 kV cm-1~6

kV cm-1, which are consistent with the non-linear out-of-plane lattice change in XRD

measurement and can be explained by electric-field-induced phase transition in

PZN-PT.

51

Figure 2.11 E-field induced FMR frequency shift under various magnetic bias fields.

The total electrically tunable range up to 13.1 GHz can be achieved with the

assistance of magnetic bias field.

By varying electric field, we have demonstrated a strong electric dependence of

FMR frequency under a certain magnetic bias field. This distinct tunability enables us

to develop a convenient engineer route to dramatically enhance tunable range and

cover whole X and K band (3 GHz~18 GHz) with great energy efficiency. Figure 2.11

shows the electric field manipulating of FMR frequency spectrum under various

external magnetic bias fields. For a small bias field 650 Oe, a large frequency tunable

range of 7.5 GHz was achieved with the minimum frequency at f=2.8 GHz. As

increasing bias fields, FMR spectrum shifts upward and reaches to the maximum

frequency of 15.9 GHz at 2060 Oe, but the frequency tunable range is slightly reduced

to 4.75 GHz. With continues shift up of FMR spectrum as increasing bias fields, a

52

strategy was proposed to dramatically enhance tunable FMR coverage with

low-power consumption. By combining two appropriate external bias fields, for

instance 650 Oe and 2060 Oe in this case, with the application of electric fields, an

enormous frequency tunable range of Δf=13.1 GHz from 2.8 GHz to 15.9 GHz was

realized, in which any FMR frequency among this range can be reached by

controlling electric fields and external magnetic bias fields. To further study how

magnetic bias field and electric field manipulate FMR frequency, theoretical

calculation based on Kittel equation as well as comparison with experimental results

are presented in Figure 2.12, where absolute value of effective magnetic field of 600

Oe is used which was derived from FMR shift in field sweeping measurement; the

magnetization of FeGaB is 9300 Gauss determined by VSM measurement. The results

show great agreement between simulation and experimental results. At low bias field,

electric field produced a large frequency tunable range of Δf with lower FMR

frequency. In contrast, at large bias field, FMR frequency was sitting at high

frequency but with less electrical tunability. From this chart, one can choose

appropriate external magnetic bias field and electric field to optimize and maximum

FMR tunability to satisfy the specific demands in real applications. So far, we have

systematically demonstrated voltage induced remarkable FMR field and frequency

tuning and proposed a technique route to dramatically enhance such tunability by

several times through combining external magnetic field and electric field. Next, we

will demonstrate a voltage pulse induced memory type of magnetization or FMR

switching, originating from the electric-field-induced phase transition in PZN-PT.

53

This is very import issue existed in most ME devices that it is challenge to realize

irreversible magnetization electrically switching.

Figure 2.12. Comparison of theoretical simulation (solid line) and experiment results

(symbol) of electric-field-induced FMR change under various magnetic bias fields.

2.3.3 Non-volatile switch of magnetism in FeGaB/PZNPT

heterostructure

Electric-field-induced phase transition is very prominent in PZN-PT single crystal

with the composition near morphotropic phase boundary (MPB). For example,

rhombohedral to orthorhombic phase transition is taken place in [011] oriented

PZN(6%~7%)-PT under a sufficient poling field. In opposite, the crystal reverted

54

back to a predominantly rhombohedral state as the remnants of orthorhombic phase is

electrically removed. Such electric-field-induced transition is irreversible due the

extra effort to overcome remnant states and is expected to display hysteresis type of

lattice change or strain vs. electric field. Given the strain induced magnetic anisotropy

and FMR change through ME coupling, phase transition as application of electric

field can be evidenced in FeGaB/PZN-PT by FMR and magnetization measurements

and could contribute to realize memory type of FMR and magnetization switching.

Figure 2.13 shows the hysteresis loops of FMR vs. electric field in field sweeping

(blue) with working frequency of 12 GHz and in frequency sweeping (red) with a

magnetic bias field of 50 Oe. Both of them exhibit a linear correlation at low electric

field, indicating a linear converse piezoelectric effect within rhombohedral phase. As

electric field reaches to the critical threshold of Ec1 ~ 5.8 kV cm-1

, a sudden change in

both resonance filed and frequency take place, suggesting the appearance of phase

transition with a remarkable lattice change and giant ME effect. At high field, FMR

field and frequency saturated with little strain variation. As lowering electric field

from 8 kV cm-1

, the orthorhombic phase and strain state remains fairly stable until

electric field reaches to another critical field of Ec2 ~3 kV cm-1

. Extremely large

converse magnetoelectric coupling coefficients of 3850 Oe·cm kV-1 (∆H/∆E), 3620

Oe·cm kV-1

(∆H/∆E) at phase transition points of 3 kV cm-1

and 5.8 kV cm-1

are

observed respectively. Symmetric behavior was observed by applying negative E-field

from 0 kV cm-1 to -8 kV cm-1 Evidenced by a sudden change of FMR field and

frequency, PZN-PT reverted back to rhombohedral phase. Note here, the opposite

55

trend in electric field dependence of resonance field and resonance frequency

(Figure5) is consistent with FMR measurements as shown in Figure 2.10 and can be

explained by Kittel equation (1). Such hysteresis type of electric field control of strain

and magnetic states provides an opportunity to realize non-volatile FMR switching,

which is extremely important in memory type ME microwave devices.

Figure 2.13 Hysteresis loops of E-field vs. FMR frequency, measured under a bias

magnetic field of 50 Oe (red) and E-field vs. FMR field with working frequency of 12

GHz (blue).

56

Figure 2.14 (a) Magnetic hysteresis loops of FeGaB/PZN-PT heterostructure

measured at different E-field; (b) Hysteresis loop of magnetization vs. E-field of

FeGaB under a magnetic bias field of 200 Oe (c) E-field impulse induced dynamically

memory-type magnetization switching.

Due to the difficulty of dynamic measurement of FMR spectrum shift, electric field

dynamically tuning of magnetization in FeGaB/PZN-PT, which could lead to a

non-volatile switching in FMR, are demonstrated as shown in Figure 2.14.

Normalized FeGaB/PZN-PT magnetic hysteresis loops under various electric fields

(Figure 2.14a) imply that a large effective magnetic field is produced and makes

magnetization process harder as 7 kV cm-1

applied. This is quite consistent with FMR

measurement. With the application of external magnetic bias field of 200 Oe, the

57

magnetization response to electric field was studied, showing an irreversible,

hysteresis type behavior as illustrated in Fig 2.14 (b), agreeing with the electric field

assisted phase transition and induced FMR hysteresis loops. Furthermore,

memory-type dynamically switching of magnetization between two bistable states

was demonstrated as shown in Figure 2.14 (c). An electric field of 5 kV cm-1

was

maintained as a bias field and 3 kV cm-1

and 7 kV cm-1

electric impulses (

58

Chapter 3 Low Temperature Fabricated Multiferroics

Heterostructure

3.1 Spin Spray deposited ZnO and Al-doped ZnO thin film

Zinc oxide is a direct wide band gap (Eg ~3.3 eV at 300 K) [35], hexagonal

wurtzite structure semiconductor that has gained a lot of attention due to its high

electrical and optical properties. It has been used in a wide variety of electronic,

optoelectronic, spintronics and nanodevices, such as, transparent thin-film transistors,

transparent electrodes in flat-panel displays and solar cells [35-40]. Compared with

other wide band gap materials, for example, GaN (Eg ~3.4 eV at 300 K), ZnO has a

large exciton binding energy (~60 meV), is more stable at high temperature, less toxic

and easier to pattern on devices. ZnO also has a much simpler crystal-growth

technology, resulting in potential low-cost ZnO-based devices.

There exists several growth methods for ZnO micro/nano-structures including

electrochemical [41],

chemical bath methods [42], hydrothermal growth [43],

chemical vapor transport [44], vapor-phase growth [45], pulsed laser deposition [46],

sputtering [47] etc. Though a lot of these conventional methods are low-temperature

and cost effective, they have the disadvantages of slow deposition rates (~0.1 nm/s)

which are too slow for industrial use. Spin-spray technique, which was original

developed for depositing high crystalline quality ferrite films from aqueous solution at

a low temperature of 90C [48], is currently been explored as a good candidate for

ZnO film growth [49-50]. The spin-spray technique is a low-temperature, low cost,

59

direct deposition technique, with the added advantage that it requires no seed layer for

film growth, has a continuous supply of fresh solution which preserves the high

concentrate of solute, and has a high deposition rate of up to 333 nm/min, making it

suitable for thick film development for industrial use. Spin-spray process has been

shown to produce high crystal quality ferrite films with good adhesion properties,

excellent magnetic propertiesand have been used for microwave devices such as

antennas and filters.

The structural, electrical and optical properties of ZnO micro/nanostructures

depend on the deposition parameters, regardless of the method of growth [51]. Wagata

et al [52], showed that the (002) peak in the XRD pattern was weakened with a

change from spin-sprayed ZnO rod array to dense film and that post annealing

affected the UV luminescence of the ZnO microstructures. Since the practical

applications of ZnO depend on these properties, it is extremely important to

investigate the effect of the deposition conditions on the growth and properties of

ZnO structures by the spin-spray technique.

To further improve the properties of the ZnO films, various elements can be doped

into their microstructures to modify their properties. By doping group-III elements

such as aluminum, boron, gallium, higher conductivity of these films can be achieved

due to oxygen vacancies and zinc interstitials [51-53]. Doped ZnO thin films also

have high temperature stability because the dopin