Rapporteurs -M Pascal

Nathalie Mme

Lemee

Maî M Freddy Ponchel

M Denis

Professeur Membres Invités -Wang Genshui

Keywords: ferroelectric, antiferroelectric, ferromagnetic, magnetoelectric composites, converse magnetoelectric effect, strain, magnetic anisotropy Ré sumé ferroé lectricité, anti-ferroé lectricité, ferromagné tisme, multiferroï ques extrinsè ques, couplage magné to-é lectrique inverse, dé formation

Growth and complete characterizations of artificial ferroics structures based on FE and AFE materials:

Performances comparison in terms of CME coupling and permeability tunability.

Croissance et caractérisations complètes de structures ferroics artificielles à base de matériaux

Ferroélectrique et Anti ferroélectrique : comparaison des performances en termes de coefficient de couplage magnétoélectrique et de l'accordabilité de la perméabilité.

JURY

Ré miens who gave this opportunity to be part of his research group. He spent a lot of time discussing the experiment results with me, helping me to find solutions for kinds of problems. I am really grateful for his valuable time and efforts on the discussion and comments of my thesis. At mean time, he always gives me kind help in my daily life of France.

Special thanks to Prof. Freddy Ponchel who gave me lots of instructions and ideas of my work. I really enjoy the time working with him for his creative thinking in research. I cannot fulfill my thesis without the help of Freddy. I will always be benefit from the impressive discussions we have ever had. I would also like to thank my colleagues from MAMINA group of IEMN-DOAE with whom I have worked closely over the years for providing me a fantastic experimental conditions and research environment. I would like to thank everyone in this warm big family for your memorable accompanies and firm support. I'd also like to especially thank Dr. Ahmad R. Hamieh for his support and help on my research and life in Valenciennes.

Furthermore, I am grateful to Prof. Tuami Lasri of IEMN for his invaluable inputs to my work. He gave me many important comments to my work, which has highly Finally, I would like to extend my gratitude to all my friends here in the France who have made my stay here so enjoyable and memorable, I would above all like to express my deepest thanks to my wonderful family for their supporting and understanding in my PhD study period. I would not be able to achieve what I have now without their supports.

Table of Contents

Motivation

Multiferroic materials, with two or more ferroic (ferroelectric, ferro/ferrimagnetic, ferroelastic) phases, [1][2][3] have offered greater potential and opportunity for developing novel multifunctional devices. Notably, the use of an electric field (E) to control the magnetic properties in multiferroic materials, known as converse magnetoelectric effect (CME), [4] has drawn intensive interest due to their essential potential in future applications such as memory devices, sensors, phase shifters, transducers, filters, oscillators, etc. [5][6][7][8][9][START_REF] Israel | A one-cent room-temperature magnetoelectric sensor[END_REF] Since single-phase magnetoelectric (ME) materials (or intrinsic multiferroic compounds, such as BiFeO3 and BiMnO3) are rare, and their ME effect is weak and usually far below room temperature, [START_REF] Ederer | Weak ferromagnetism and magnetoelectric coupling in bismuth ferrite[END_REF][START_REF] Lawes | Introduction to magnetoelectric coupling and multiferroic films[END_REF] most researchers have focused on ME composites to achieve strong ME coupling effect. [START_REF] Catalan | Physics and applications of bismuth ferrite[END_REF][START_REF] Srinivasan | Magnetoelectric composites[END_REF][START_REF] Vaz | Magnetoelectric coupling effects in multiferroic complex oxide composite structures[END_REF] The ME coupling effect in composites or heterostructures can be exploited by the strain transferring between the ferroelectric (FE) phases and ferromagnetic (FM) phases.

Much advancement has recently been achieved in FM/FE composites, such as

CoFe2O4/BaTiO3, Ni0.8Zn0.2Fe2O4/Pb(Zr, Ti)O3, Terfenol-D/Pb(Mg, Nb)O3-PbTiO3, FeBSiC/PVDF. [START_REF] Islam | Effect of piezoelectric grain size on magnetoelectric coefficient of Pb(Zr0.52Ti0.48)O3-Ni0.8Zn0.2Fe2O4 particulate composites[END_REF][START_REF] Fang | Enhancing the magnetoelectric response of Metglas/polyvinylidene fluoride laminates by exploiting the flux concentration effect[END_REF][START_REF] Vopson | Fundamentals of Multiferroic Materials and Their Possible Applications[END_REF][START_REF] Erdem | Nanoparticle-Based Magnetoelectric BaTiO 3 -CoFe 2 O 4 Thin Film Heterostructures for Voltage Control of Magnetism[END_REF] An application of the external electric field can lead to large magnetization and susceptibility changes in these systems. Nevertheless, investigations for a better understanding of fundamental physics and engineering potential of these materials and structures are still on the way.

Previous researchers have studied a variety of ME composites. [7,[START_REF] Liu | Voltage tuning of ferromagnetic resonance with bistable magnetization switching in energyefficient magnetoelectric composites[END_REF][START_REF] Wu | Recent progress on the structural characterizations of domain structures in ferroic and multiferroic perovskite oxides: A review[END_REF][START_REF] Mirshekarloo | Ferroelastic strain induced antiferroelectric-ferroelectric phase transformation in multilayer thin film structures[END_REF] The main interest has been focused on FM/FE heterostructures and in particular the tuning of the magnetic response through the electric-field-induced strain of the FE layer applied to the FM layer. As antiferroelectric (AFE) materials enjoy a large strain during the electric-field induced AFE-FE transition, and display zero polarization and strain in a relaxed state, [START_REF] Lepadatu | Heat-assisted multiferroic solid-state memory[END_REF][START_REF] Hao | A comprehensive review on the progress of lead zirconate-based antiferroelectric materials[END_REF] there is a potential benefit for their use. Thus the advantages of large E-induced phase transition strain and zero remnant state in AFE materials, [START_REF] Zhou | Antiferroelectric materials, applications and recent progress on multiferroic heterostructures[END_REF] make that AFE-based ME composites are also interesting. However, the reported works of electric-field control of magnetism in AFE-based ME composites are scarce when compared to the work on FM/FE composites. [START_REF] Zhou | Strong non-volatile voltage control of magnetism in magnetic/antiferroelectric magnetoelectric heterostructures[END_REF][START_REF] Zhou | Magnetoelectric coupling in antiferroelectric and magnetic laminate composites[END_REF] Thus, comprehensive research of E-field controlled magnetism in AFE-based ME heterostructures is necessary to understand the fundamental of ME coupling effect in FM/AFE heterostructures. At the same time, to reveal the differences of ME coupling effect in FM/AFE and FM/FE heterostructures, the comparison of E-field controlled magnetism in AFE-and FEbased ME heterostructures is also important.

In past decades, many different ferromagnetic or ferrimagnetic materials are used as magnetic phases to develop ME composites when targeting different applications, including oxides such as spinel ferrites, hexagonal ferrites, and lanthanum manganites, 3d-transition metals Fe, Co, and Ni, alloys such as permendur, Terfenol-D, Galfenol. [9,[START_REF] Petrov | Dispersion characteristics for low-frequency magnetoelectric coefficients in bulk ferritepiezoelectric composites[END_REF][START_REF] Bai | Magnetic and magnetoelectric properties of as-deposited and annealed BaTiO3-CoFe2O4 nanocomposite thin films[END_REF][START_REF] Zhang | Enhancement in magnetoelectric response in CoFe2O4-BaTiO3 heterostructure[END_REF][START_REF] Li | Novel NiZnAl-ferrites and strong magnetoelectric coupling in NiZnAl-ferrite/PZT multiferroic heterostructures[END_REF][START_REF] Yao | Giant magnetodielectric effect in Terfenol-D/PZT magnetoelectric laminate composite[END_REF][START_REF] Shen | High non-linear magnetoelectric coefficient in Metglas/PMN-PT laminate composites under zero direct current magnetic bias[END_REF][START_REF] More-Chevalier | Reversible control of magnetic domains in a Tb0.3Dy0.7Fe2/Pt/PbZr0.56Ti0.44O3 thin film heterostructure deposited on Pt/TiO2/SiO2/Si substrate[END_REF][START_REF] Dusch | Stress-mediated magnetoelectric memory effect with uni-axial TbCo2/FeCo multilayer on 011-cut PMN-PT ferroelectric relaxor[END_REF] In strain-mediated CME heterostructures, a strong coupling between elastic and magnetic domains is required for the magnetic component. [START_REF] Wan | Electric-field-induced magnetization in Pb(Zr,Ti)O3/Terfenol-D composite structures[END_REF] The ferromagnetic materials, such as Ni-Mn-Ga (NMG) alloys and Tb, Co, Fe metals, have attracted much attention as they combine both magnetic and elastic ordering and present high magnetostriction coefficient. [START_REF] Dusch | Stress-mediated magnetoelectric memory effect with uni-axial TbCo2/FeCo multilayer on 011-cut PMN-PT ferroelectric relaxor[END_REF][START_REF] Wang | Coupling of ferroelastic strain and ferroelectric phase transition in NiMnGa/Pb0.97La0.02(Zr0.95Ti0.05)O3 bilayered films[END_REF][START_REF] Tiercelin | Enhanced magnetoelectric effect in nanostructured magnetostrictive thin film resonant actuator with field induced spin reorientation transition[END_REF][START_REF] Klimov | Magnetoelectric write and read operations in a stress-mediated multiferroic memory cell[END_REF] These properties make it possible to easily manipulate their ferroelastic domains by applying a magnetic field or mechanical stress. [START_REF] Li | Stress-induced change of magnetization in a Ni-Mn-Ga single crystal under magnetomechanical training[END_REF] These characteristics provide opportunities to induce a substantial magnetization change by external mechanical stimulus and reach a high ME coupling coefficient, which is suitable for strain-mediated CME coupling. Ferrite yttrium iron garnet (YIG) is also considered to be a good choice as the ferrimagnetic material to develop ME composites when targeting tunable multiferroic microwave devices thanks to its outstanding advantages including small magnetocrystalline anisotropy, narrow ferromagnetic resonance linewidth, high resistivity, and low microwave losses.

Thus three magnetic materials (ferrimagnetic oxide YIG film, the ferromagnetic NMG alloys and TbCo2/FeCo multilayers) are used to couple with FE and AFE materials and to compare the performance of E-field controlled magnetism in AFE-and FE-based ME composites. In this thesis, we focused on the CME coupling effect in both FE-and AFE-based ME composites. In order to observe the distinct coupling effect, the composites' laminated structure is designed to be FM films deposited on FE or AFE ceramic substrates. The voltage-controlled magnetization switching and magnetic susceptibility variations are studied in different heterostructures.

Scope of Work

The thesis begins with a background review in Chapter 2. The basic theory and general concepts of multiferroic and ME materials would be introduced. The research progress on AFE-based multiferroic material fabrication, mechanism studies, and practical applications are given in the past few decades, as well as the current problems needed to be solved and identified.

In Chapter 3, the main experimental techniques would be given. It includes ceramic and thin film preparation, structural analysis, electrical and magnetic characterization.

Backgrounds and Literature Review

The concept of "ferroic" is referred to the properties of ferroelectricity, ferromagnetism, and ferroelasticity in a crystal, as Aizu concluded in 1970. [START_REF] Aizu | Possible species of ferromagnetic, ferroelectric, and ferroelastic crystals[END_REF] The most distinct feature of "ferroic" is the existence of domain. The order parameters (dipole moment, magnetic moment, etc.) of atoms or ions in adjacent positions of ferroic materials are arranged in the same direction in the local area due to interaction, thereby forming an ordered region, named "domain". The domains allow the order parameter of ferroic materials to change with the presence of the external field, and leads to hysteresis loops. Later, it was found that ferrotoroidicity also belongs to the family of "ferroic". [START_REF] Schmid | Some symmetry aspects of ferroics and single phase multiferroics*[END_REF] The spontaneous polarized or magnetized domains in ferroic materials allow the order parameter to change with the application of the external fields, and ferroic materials would present hysteresis loops and anomalies at the phase transition.

Theoretical Background

Ferroelectrics and Antiferroelectrics

The study of ferroelectric materials is generally believed to have begun in 1921, when J. Valasek firstly observed the dielectric anomalies in Rochelle salt. [START_REF] Valasek | Piezo-electric and allied phenomena in rochelle salt[END_REF] Then, ferroelectric materials develop quickly and become essential components in a broad spectrum of applications. The concept of antiferroelectricity was proposed in 1951 to describe a state where chains of ions in the crystal are spontaneously polarized, but with neighboring chains polarized in antiparallel directions. [START_REF] Kittel | Theory of antiferroelectric crystals[END_REF] Since then, various antiferroelectrics have been discovered over the past half century, and research on antiferroelectric materials has received increasing attention.

(1) Ferroelectrics Ferroelectric materials are a family of materials who display spontaneous polarization in a specific range of temperature, and the polarization of FE materials is reversible with an application of the electric field. In FE materials, the adjacent dipoles in one domain share the same polarization orientation, and orientation of the dipoles can be aligned by an external DC electric field. For illustration, Figure 2.2 (a) shows the electric-field dependence of polarization (P-E) of FEs. When the external electric field turns to the opposite direction, new ferroelectric domains of FEs begin to nucleate and grow up, the domain wall moves, and finally, the polarization reverses. When an electric field is applied to ferroelectric materials, a strain accompanying the polarization reversal process would also be generated in ferroelectrics, as given in Fig. 2.2 (b). The electric-field dependence of strain (S-E) loop in a butterfly shape can be attributed to three effects: the converse piezoelectric effect, the switching, and movements of domain walls. The hysteresis loop of FEs indicates that the spontaneous polarization is reversible under an application of the electric field. This feature is essential for the applications of data storage, especially non-volatile memory devices. In addition to ferroelectricity, ferroelectric materials also present dielectric, piezoelectric, pyroelectric properties, which shows great potential in the applications of semiconductor and applied physics.

(2) Antiferroelectrics Although the precise definition of antiferroelectrics is still controversial, it is generally believed that the adjacent dipoles are aligned in opposite orientation in AFE materials. In this case, if a sufficiently high DC electric field is applied, the polarization orientation of dipoles would be aligned into one direction so as to realize an FE state. This is called the electric-field-induced phase transition.

When an electric-field induced AFE-FE phase transition happened, the corresponding structure phase transition with an apparent expansion of volume would generate big cell strain, as given in Fig. 2.3 (b). Unlike FE materials, the strain or displacement induced in AFE materials originates from the E-induced AFE-FE phase transition because the unit cell of the FE phase is larger than that of the AFE phase.

Hence it is possible to keep the high strain level even when the materials are in thinfilm form. Figure 2.4 (a) shows the double P-E hysteresis loops of AFE materials. [START_REF] Hao | A comprehensive review on the progress of lead zirconate-based antiferroelectric materials[END_REF] When the applied electric field is lower than the phase transition field strength (abbreviated as EF), the AFE materials exhibit the properties of a linear dielectric. When the applied electric field is higher than EF, the polarization jumps suddenly to a very high level and is close to saturation. The field-induced FE phase returns to the AFE phase when the electric field is removed, or the electric field is reduced to a field-induced FE phase to a critical electric field (EA) of the AFE phase. This reversible process, accompanied by a significant strain reduction because the unit cell shrinks exhibits a sizeable electrostrictive coefficient. Therefore, the S-E hysteresis loop of AFEs is entirely different from the S-E loop of FEs, as shown in Fig. 2.4 (b). The novel E-induced phase transition of AFEs has drawn tremendous interest for their engineering applications, such as high energy storage capacitors, micro-actuators, pyroelectric security sensors, cooling devices, and pulsed power generators and so on. [START_REF] Zhao | Energy-Storage Properties and Electrocaloric Effect of Pb(1-3x/2)LaxZr0.85Ti0.15O3 Antiferroelectric Thick Films[END_REF][START_REF] Ahn | Antiferroelectric Thin-Film Capacitors with High Energy-Storage Densities, Low Energy Losses, and Fast Discharge Times[END_REF][START_REF] Li | High-field antiferroelectric behaviour and minimized energy loss in poly(vinylidene-co-trifluoroethylene)graft-poly(ethyl methacrylate) for energy storage application[END_REF][START_REF] Chauhan | Anti-Ferroelectric Ceramics for High[END_REF] In recent years, some researchers turned the attention to AFE-based multiferroic composites because of their considerable electric-field induced strain. [START_REF] Zhou | Antiferroelectric materials, applications and recent progress on multiferroic heterostructures[END_REF]

Magnetism and Magnetic Properties

The first deliberate use of the magnetism stems from the 11th century in forms of the magnetic compass, which has dramatically encouraged human navigation and geographic discoveries. The remarkable discovery of Hans Christian Oersted in 1821 that an electric current flowing through a wire produces a magnetic field, has opened a new page of fundamental magnetic research. [START_REF] Cullity | Introduction to magnetic materials[END_REF] Since then, significant development of the electromagnets has been achieved, which made major contributions to the development of the functional materials.

(1) Magnetism Magnetism is a quantum mechanical phenomenon originating from the fundamental properties of the atom and its electrons. There are two types of rotational motions for electrons. One is called "spin" that the electron spins around itself, another one is "orbit" that the electron moves along an orbit around the atom. The electrons orbital motion and spin themselves can create magnetic moments, respectively, and they also can be coupled, which is called spin-orbit coupling. [START_REF] Spaldin | Magnetic materials: fundamentals and applications[END_REF] The coupling also includes orbit-orbit coupling and spin-spin coupling among electrons, which give a total atomic magnetic moment. Thus all materials experience magnetism, but in practice, not all materials can produce effects that are reasonably large.

The magnetic behavior of different materials under the influence of a magnetic field is given as

𝑀 = 𝜒𝐻 (2.1)
where M is the magnetization, χ is the magnetic susceptibility, and H is the magnetic field strength. There are five types of magnetic materials concerning their magnetic characteristics. They are diamagnetic, paramagnetic, antiferromagnetic, ferromagnetic, and ferrimagnetic materials, as given in Fig. 2.5.

Diamagnetic, paramagnetic and antiferromagnetic materials are weak magnetic materials, while ferromagnetic and ferrimagnetic materials are of strong magnetism.

Diamagnetic materials have negative and small susceptibilities (χ= 10 -5 ), hence repulse magnetic fields. Paramagnetic materials exhibit positive but small susceptibilities in the range of χ = 10 -5 -10 -2 . Antiferromagnetic materials magnetize weakly in the direction of an external magnetic field. In antiferromagnetic materials, the adjacent magnetic moments are aligned in opposing direction with the same magnitude, like antiferroelectrics. In ferrimagnetic materials, the adjacent magnetic moments are also aligned in opposing direction, but with different magnitudes, thus magnetic susceptibilities of ferrimagnetic materials are high. Ferromagnetic materials retain almost perfectly aligned magnetic moments even without an external magnetic field that they are the most widely utilized magnets with very high susceptibilities. (2) Ferromagnetic and Ferrimagnetic Materials (FM Materials)

Ferromagnetic and ferrimagnetic materials present two distinct characteristics: spontaneous magnetization and magnetic domain. The strong internal interactions of ferromagnetic materials align the magnetic moments parallel to each other, which produces a spontaneous magnetization even in the absence of an external magnetic field. [START_REF] Morrish | The physical principles of magnetism[END_REF] The spontaneous magnetization orients to preferred directions because of the magnetocrystalline anisotropy. Ferromagnetic domains are small regions where all the atomic magnetic dipoles have a parallel alignment to each other. [START_REF] Hubert | Magnetic domains: the analysis of magnetic microstructures[END_REF] The magnetization directions of different domains can be different. Likewise, ferrimagnetic materials have populations of atoms with opposing magnetic moments, as in antiferromagnetism;

however, in ferrimagnetic materials, the opposing moments are unequal. Thus the spontaneous magnetization and magnetic domains remain. With an application of the external magnetic field, the domain wall motion and domain rotation work together in ferromagnetic materials and lead to a magnetizationmagnetic field (M-H) hysteresis loop, as given in Fig. 2.6. In ferromagnetic materials, all domains are randomly arranged, and the total magnetization is zero with the absence of an external magnetic field. When the external magnetic field is applied, the domain in which the spontaneous magnetization is nearest to the applied field direction starts to grow with domain wall moving. The other domains whose spontaneous magnetization is not close to the applied field direction, become smaller and smaller.

Finally, all the domains align with the applied field direction; hence, the magnetization reaches saturation, as illustrated in Fig. 2.7.

Figure 2.7

The magnetic domain of ferromagnetic materials.

These features of ferromagnetic materials provide great opportunities in a wide range of applications in modern device technologies.

Magnetoelectric materials

(1) Multiferroics

In the past decades, multiferroics have attracted many interests due to its significant improvement to data storage, sensors, and spintronics devices. In the definition, multiferroic materials have any two of the above ferroic orders and interactions, as shown in Fig. 2.8. The ME response is the appearance of an electric polarization (P) upon applying a magnetic field (H) (i.e., so-called direct ME effect):

∆𝑃 = 𝛼 • ∆𝐻 𝑜𝑟 ∆𝐸 = 𝛼 𝐸 • ∆𝐻 (2.2)
where E denotes electric field, and α (αE) is the ME (ME voltage) coefficient. Thus, the electric polarization can be modified by a magnetic field, or voltage output can be produced by applying a magnetic field on the materials via the ME coupling.

Sometimes, the magnetodielectric effect is used, i.e., the change in dielectric constant with the magnetic field. In addition to this direct ME effect, there is simultaneously the converse ME (CME) effect, i.e., the appearance of magnetization (M) upon applying an electric field (E),

∆𝑀 = 𝛼 • ∆𝐸 (2.3)
which indicates the modulation of magnetism by an electric field. In a multiferroic material with ferroelectric and ferromagnetic orders co-existing, such strong ME effects make the multiferroic ME materials promising in applications such as multifunctional sensors, transducer, filters, oscillators, phase shifters, memory devices, and so on.

(

2) Classifications of ME Materials

There are two categories of ME materials, single-phase ME materials and composites.

Single-phase ME materials exhibit ferroelectricity and ferromagnetism in one single compound. The general criteria for the appearance of ferromagnetism is exchanging interaction of electron spins, while the redistribution of charge density in the lattice for the ferroelectricity. These make it possible for the coexistence of spontaneous magnetization and polarization in one single phase ME material. [START_REF] Hill | Why are there so few magnetic ferroelectrics?[END_REF] Single-phase ME materials are chemically homogenous compounds, and they exhibit intrinsic ME coupling by the co-existence of magnetic moments and electric dipoles simultaneously. In 1957, scientists predicted the existence of a linear ME coupling coefficient in Cr2O3; subsequently, the ME coupling effect was experimentally confirmed with a ME coefficient of α = 4.1 pS/m. [START_REF] Astrov | The magnetoelectric effect in antiferromagnetics[END_REF][START_REF] Folen | Anisotropy of the magnetoelectric effect in Cr2O3[END_REF] Since then, various singlephase ME materials (GaFeO3, Ti2O3, REMnO3, RE = Y, Ho, Er, Tm, Yb, Lu, etc.) have been gradually discovered, [4] the representative of single-phase ME materials is BiFeO3 (BFO). BFO is a room-temperature single-phase multiferroic materials with the coexistence of robust ferroelectricity and antiferromagnetism (TFE = 1100 K, TN = 643K), and a weak canted magnetic moment at room temperature. It is predicted that the saturated polarization of BFO can reach 100 μC/cm 2 at room temperature. [START_REF] Zhao | Electrical control of antiferromagnetic domains in multiferroic BiFeO3 films at room temperature[END_REF] However, the ME coupling effect is restricted in BFO because of its weak ferromagnetism above room temperature. [8] In general, the current ME coupling effect of most single-phase ME materials are still weak, which greatly hinders their applications. Therefore, the researchers turned their attention to ME composites, which consist of physically separated magnetic and electric order phases. By artificially design of ferromagnetic/ferroelectric multiphase composites, the ferroelectric phase, and ferromagnetic phase can be selected from a wide range of materials to obtain the optimized ME coupling effect, which offers a larger space to design and manipulate ME materials, and they exhibit higher ME coupling coefficient than those found in single-phase materials at room temperature.

For example, it is several orders of higher magnitude than that in single-phase multiferroic materials so far available at room temperature. [START_REF] Nan | Multiferroic magnetoelectric composites: Historical perspective, status, and future directions[END_REF] The concept of ME composites was first proposed in 1972, [START_REF] Van Suchtelen | Product properties: A new application of composite materials[END_REF] then various multiferroic composites have been reported. [START_REF] Devan | Electrical conduction and magnetoelectric effect of (x) BaTiO3 + (1-x) Ni0.92Co0.03Cu0.05Fe2O4 composites in ferroelectric rich region[END_REF][START_REF] Zhou | Dielectric, magnetic, and magnetoelectric properties of laminated PbZr0.52Ti0.48O3∕CoFe2O4 composite ceramics[END_REF][START_REF] Devan | Ferroelectric and ferromagnetic properties of (x)BaTiO3 -(1-x)Ni0.94Co0.01Cu0.05Fe2O4 composite[END_REF][START_REF] Li | Enhanced magnetoelectric effects in composite of piezoelectric ceramics, rare-earth iron alloys, and ultrasonic horn[END_REF][START_REF] Zavaliche | Electric field-induced magnetization switching in epitaxial columnar nanostructures[END_REF][START_REF] Lei | Straincontrolled magnetic domain wall propagation in hybrid piezoelectric/ferromagnetic structures[END_REF][START_REF] Lei | Electric field control of domain wall logic in piezoelec-tric/ferromagnetic nanodevices[END_REF] The phase connectivity and interfacial bonding of the FE and FM phases can greatly influence the ME coupling effect. Based on the phase connectivity of composites, [START_REF] Newnham | Connectivity and piezoelectricpyroelectric composites[END_REF] ME composites can be commonly prepared through three different ways: 0-3, 2-2, and 1-3 connectivity, as shown in Fig. 2.9 (a), (b) and (c). The 0-3 particle-matrix composites mainly combine ferroelectric or magnetic oxide particles into another phase matrix for sintering, thereby achieving two-phase composites. The 0-3 type is simple in preparation, and the interface between two phases is large enough for strain transfer, but the magnetic particles with good conductivity would lead to high leakage, which hinders its development. (3) The Strain Mediated ME Coupling Effect

The ME coupling effect in ME composites can be realized by three different mechanisms, i.e., charge mediated ME effect, [START_REF] Molegraaf | Magnetoelectric effects in complex oxides with competing[END_REF][72][73] exchange bias mediated ME effect and strain mediated ME effect. [START_REF] Lei | Straincontrolled magnetic domain wall propagation in hybrid piezoelectric/ferromagnetic structures[END_REF][START_REF] Chu | Electric-field control of local ferromagnetism using a magnetoelectric multiferroic[END_REF][START_REF] Yu | Interface ferromagnetism and orbital reconstruction in BiFeO3-La0.7Sr0.3MnO3 heterostructures[END_REF][START_REF] Zavaliche | Electrically assisted magnetic recording in multiferroic nanostructures[END_REF][START_REF] More-Chevalier | Magnetoelectric coupling in Pb(Zr,Ti)O3-Galfenol thin film heterostructures[END_REF] The strain mediated ME effect in composites results from the elastic coupling between the piezoelectric and magnetostrictive components.

The strain mediated mechanism is to manipulate the ferromagnetic or ferroelectric response with the electric field or magnetic field induced strain, as given in Fig. 2.10.

In direct ME (DME) coupling, the applied magnetic field generates strain in the magnetic layer via the magnetostriction effect, and this strain is transferred to the piezoelectric layer resulting in an electric displacement or a dielectric polarization through the piezoelectric effect. In CME coupling, an external electric field induces strain in the ferroelectric layer due to the inverse piezoelectric effect, and the strain transferred to the magnetic layer produces a magnetization change or domain reorientation by magnetic anisotropy transformation. In a word, the strain can be transferred through the interface from one phase to another phase, thereby realizing the H-field control of polarization or E-field manipulation of magnetization. The strain mechanism is the most common and effective method to realize the ME coupling effect in the ME composites due to its simple principle and strong practicability.

In this strain-mediated ME coupling effect, the interface between FE phases and FM phases plays a crucial role in ME response. [START_REF] Liu | Dependence of giant magnetoelectric effect on interfacial bonding for multiferroic laminated composites of rare-earth-iron alloys and lead-zirconate-titanate[END_REF][START_REF] Nan | Influence of interfacial bonding on giant magnetoelectric response of multiferroic laminated composites of Tb1-xDyxFe2 and PbZrxTi1-xO3[END_REF][START_REF] Zhou | Effects of the volume fraction of piezoelectric particles in the magneto-electro-elastic interfacial region on the fracture behavior of a laminate multiferroic plate[END_REF][START_REF] Chen | Bending of multiferroic laminated rectangular plates with imperfect interlaminar bonding[END_REF] Several researchers have demonstrated that the large thickness and low shear modulus of the interfacial bonding results in a significant drop in ME response.

(

4) The Strain-Induced Magnetic Anisotropy Transformation

To reveal how the strain mediated CME coupling effect works, it is necessary to demonstrate the strain-induced magnetic anisotropy transformation in magnetic materials. The strain can induce a magnetic anisotropy in magnetic materials due to the magnetoelastic effect. For example, in polycrystalline magnetic films, there is not a preferred orientation due to the random grain arrangement. Thus the magnetocrystalline anisotropy energy is small. The effect of elastic stress on magnetization is called the inverse magnetostrictive effect. The applied elastic stress will alter the domain structure and create a new anisotropy competing with the other anisotropies. The effect of elastic stress on magnetization can be understood based on minimizing magnetoelastic energy in the magnetic system. The magnetoelastic energy (Eme) can be given by [START_REF] Cullity | Introduction to magnetic materials[END_REF] 𝐸 𝑚𝑒 = -

3 2 𝜆 𝑠 𝜎𝑐𝑜𝑠 2 𝜃 (2.4)
where λS is the saturation magnetostriction constant, σ is the external stress, and θ is the angle between the saturation magnetization (MS) and the stress (σ).

If the magnetostriction constant λ is positive in a material, the external tensile stress (σ > 0) will lead more magnetic moments to align along the stress direction (i.e., θ = 0°) to reach the minimum of magnetoelastic energy. When a magnetic field is applied along this tensile stress direction, a lower magnetic field will be required to saturate the sample, which means that tensile stress produces a preferred magnetization direction. When compressive stress is applied (σ < 0) to a magnetic material, which has a positive λ, magnetic moments tend to align perpendicular to the compressive direction (i.e., θ = 90°) to reach the minimum of magnetoelastic energy.

In a word, E control of magnetism in multiferroic heterostructures is achieved by E-created magnetic anisotropy. Strain mediated ME coupling in layered ferromagnetic/ferroelectric heterostructures can provide great opportunities in realizing novel multiferroic devices.

(5) Measurements of CME Generally, there are three ways to characterize the CME effect, i.e., the voltage induced magnetization variations, E-tuned ferromagnetic resonance (FMR) shift in the resonance frequency and in situ E-field controlled magnetic domain change by scanning probe microscopy techniques.

According to Equation 2.3, the electrically induced ME coupling coefficient can be estimated by measuring the induced magnetization with an applied AC voltage. The CME coefficient can be determined by the slope of the linear graph of the induced magnetization versus the AC voltage amplitude. [START_REF] Zhang | Enhanced converse magnetoelectric effect in Pb(Zr,Ti)O3-bimorph/Metglas laminated composite[END_REF][START_REF] Vopson | Measurement techniques of the magneto-electric coupling in multiferroics[END_REF] Another commonly used CME coefficient can be deduced from the FMR shifts under an applied electric field that manifests the E-induced strain to be an internal magnetic field in composites. The effect is expected to be the largest when the frequency of the AC electric field/voltage matches the electro-mechanical resonance frequency of the sample. [START_REF] Vopson | Measurement techniques of the magneto-electric coupling in multiferroics[END_REF] The ME coupling can also be in situ characterized by scanning probe microscopy techniques, such as magnetic force microscopy (MFM), and FMR under the electric field with the use of a scanning microwave microscopy (SMM). [START_REF] Srinivasan | Magnetoelectric composites[END_REF][START_REF] Nan | Multiferroic magnetoelectric composites: Historical perspective, status, and future directions[END_REF][START_REF] Ma | Recent progress in multiferroic magnetoelectric composites: from bulk to thin films[END_REF]

Literature Review

Since AFE materials, with novel properties of electric-field-induced AFE/FE phase transition, exhibit a large and sharp strain change during the phase transition for atomic rearrangements and zero polarization and strain in relaxed state. On the one hand, the significant strain of AFEs, which can reach 0.87 % in some reported AFE ceramics, make them extremely attractive for strain-mediated ME coupling. [START_REF] Pan | Field-forced antiferroelectric-toferroelectric switching in modified lead zirconate titanate stannate ceramics[END_REF][START_REF] Sharifzadeh Mirshekarloo | Large strain and high energy storage density in orthorhombic perovskite (Pb0.97La0.02)(Zr1-x-ySnxTiy)O3 antiferroelectric thin films[END_REF][START_REF] Pan | Large displacement transducers based on electric field forced phase transitions in the tetragonal (Pb 0.97 La 0.02 ) (Ti,Zr,Sn)O 3 family of ceramics[END_REF] On the other hand, the non-remnant state of AFEs results in no pre-poling process which shows excellent compatibility with various applications. [START_REF] Zhou | Antiferroelectric materials, applications and recent progress on multiferroic heterostructures[END_REF] To explore the physics underlying the ME effects, researchers have developed state memory in an FM/AFE bilayer. [START_REF] Lepadatu | Heat-assisted multiferroic solid-state memory[END_REF] In this model, the strain is generated using the AFE layer by voltage-induced AFE-FE phase change, and it is transferred to the magnetic free layer by strain-mediated coupling effect. The thermally activated straininduced magnetization switching was analyzed using a three-dimensional, temperature- Thus the magnetic properties (magnetization and FMR field) can be tuned by voltage.

The nonvolatile voltage control of magnetism is showing promising applications in compact and power-effective multiferroic devices.

The

Experimental Techniques

In this chapter, the experimental techniques used in this thesis are presented in detail. The fabrication of polycrystalline thin films and ceramics, the structural characterization and property characterizations are discussed.

Ceramic Synthesis

The ceramics preparation was finished by the author in Shanghai Institute of Ceramics, Chinese Academy of Science. The choice of FE (PMN-PZT) and AFE (PLZST) ceramics will be explained in chapter 4.

Raw materials

The ceramics were prepared by a conventional solid-state reaction. The raw powders used in this work are shown in Table 3.1.

Solid-State Reaction Process

The raw powders were prepared by the conventional solid-state method according to the chemical formula of the ceramics with 0.5 wt.% excess Pb3O4 to compensate Pb volatilization during sintering. The raw powders were weighed according to the compositions by the balance. Then, they were mixed with deionized water and milling ball according to the weight proportion of 1: (1.4 ~1.8) : (0.7 ~1.0). After 10 hours of ball milling, the mixed powders dried in the oven, sieved with mesh and then calcined at 850 ℃ for 2 hours. After calcination, the calcined powder was milled again for 24 hours and dried overnight before mixing with 6 wt.% of polyethylene glycol as a binder.

Next, the calcined powders were uniaxially pressed into pieces of the disc with a thickness around 2 ~3 mm and a diameter of 15 mm. To decompose the binder fully, the green pellets were firstly put at 800 ℃ for 2 hours. Then for sintering, the green pellets were covered in a calcined powder of the same composition in a closed alumina crucible to avoid volatilization of Pb and sintered at 1240 ~1320 °C for 2 to 4 hours.

Note that the heating and cooling rate used for both calcination and sintering was 2 °C/min.

Processing Techniques

The sintered pellets were mechanically ground to a final uniform thickness of approximately 0.5 mm. Then the silver electrodes were coated on two sides of the clean ceramic pellets through the screen printing technique, and the silver sintering was conducted at 700 ~750 °C for 15 ~30 minutes. Finally, the ceramic samples were ready for electrical measurements.

The sintered ceramics were machined and polished on both sides into the ceramic substrate. The roughness of the surface after polishing is about 10 nm. The substrate dimension is 10 × 10 × 0.5 mm 3 .

Thin Film Preparation

Magnetron Sputtering

Sputtering is one of the most common physical vapor deposition (PVD) techniques for fabricating thin films onto substrates. The sputtered films tend to be more uniform and dense with fine grain size (~ 10 nm) compared to other PVD methods.

[100] Fig.

3

.1 shows a schematic of a magnetron sputtering system. In this system, the target materials are connected to the negative terminal of a power supply, typically about a few kilovolts (kV), while the substrate may be grounded or biased positively. [START_REF] Ohring | Materials science of thin films[END_REF] The 29 chamber is first pumped down to a vacuum with the base pressure about 10 -5 Pa, and then the working gas (typically, argon) is introduced and serves as a visible glow discharge medium between two electrodes. In the discharge, positive gas ions and electrons are produced and form a plasma. The ions strike the target and physically remove the target atoms. The target atoms enter and pass through the plasma region to eventually deposit onto the substrate and to form the film. In this work, the Plassys RF magnetron sputtering system was chosen to deposit YIG and NMG thin films onto Si substrate and polished ceramic substrates, as given in Fig. 3.2.

Targets Fabrication

Before the sputtering, the targets should be prepared first. In this work, the raw powders of Y2O3 and Fe2O3 were used to fabricate the YIG target. The raw powders were weighed according to the compositions by the balance. The total weight of raw powders was about 60 g. Then, they were mixed with propan-2-ol and milling ball for 10 hours of ball milling. Next, the mixed powders dried in the oven, sieved with mesh.

Finally, the powders were pressed in cylindrical steel under the uniaxial pressure of 5, 10, 15, 20 Tons for 5 minutes, respectively. The targets are of thickness around 3 ~4 mm and a diameter of 3 inches.

The NMG magnetic alloy target with a composition of Ni0.5Mn0.33Ga0.17 was bought from the company.

Electrodes Deposition

The Pt electrodes were included in some heterostructures, and metallic masks were used to confine the dimension of the electrodes. The Pt electrodes were deposited by DC magnetron sputtering process on Si or ceramic substrates.

Characterization Methods

Microstructural Characterization

(1) X-Ray Diffraction (XRD)

X-ray diffraction is powerful to determine the crystal structure, symmetry, and lattice parameters. In this study, we utilize an X-ray diffractometer (Rigaku D/MAX-2550, Japan) to characterize the phase structure of ceramics and thin films. The X-ray tube was operated at 40 kV and 40 mA with a wavelength of λ=0.154nm (Cu Kα).

(2) Atomic Force Microscope (AFM)

Atomic force microscope is to characterize the morphology, roughness, and other properties of surfaces on the atomic scale at environmental conditions. In this work, the surface morphology and roughness of the samples were measured using non-contact AFM (Nanonavi SII, Japan).

(3) Scanning Electron Microscope (SEM)

The scanning electron microscope is a useful magnification tool that utilizes focused beams of the electron to obtain information and provides high-resolution, threedimensional images. The field emission SEM (JEOL/JSM-6700F, HITACHI/ FESEM-S4800, Japan) is used to characterize the morphology of thin films.

(4) Magnetic Force Microscope (MFM)

Magnetic force microscopy (MFM) is a kind of AFM where a magnetized tip can scan the magnetic sample and characterize the magnetic structure of the sample surface.

In this work, the MFM (MultiMode, Bruker) measurements were performed with an external voltage applied to the sample by two probes at the same time to characterize the magnetic domain change with the electric field in ME composite films.

Electrical Measurements

The metal-insulator-metal (MIM) capacitor structure is built by depositing electrodes on the samples' surfaces.

(1) Dielectric Properties

The dielectric constant was measured using an LCR meter at a frequency of 100 -100k Hz. The impendence analyzer HP 4284 and HP 4294A connected with VT 7004 and GJW-1 furnace. By measuring the capacitance of a sample, the dielectric constant is obtained from an equation as

ε 𝑟 = 𝐶𝑑 𝐴𝜀 0 (3.1)
where C is capacitance (F), 𝜀 0 is vacuum permittivity (8.85×10 -12 F/m), A is the area of electrode surface (m 2 ), and d is the thickness of the sample (m).

(

2) Ferroelectric Properties

The polarization-electric-field (P-E) hysteresis loops of samples were measured by TF Analyzer 2000 (aixACCT) system. The P-E loops can be measured in a temperature range from -40 ℃ to 200 ℃ and with a frequency of 1-5000 Hz.

(

3) Strain Measurements

The strain-electric-field (S-E) curves of samples were also measured by TF Analyzer 2000 (aixACCT) system. The ceramics were characterized by TF Analyzer 2000 equipped with a laser interferometer vibrator.

Magnetic Properties Characterization

The vibrating sample magnetometer (VSM) system is used to test the lowfrequency magnetic properties of magnetic materials. A typical VSM system includes an electromagnet to provide the DC magnetic field, a vibrating stage to vibrate the sample, and a pick-up coil assembly that measures the magnetization of the sample according to Faraday's law of induction. The output is usually the magnetic moment M or flux B vs. the field H. The magnetic hysteresis loops (M-H) are measured at room temperature by using a VSM (ADE model EV9) with a vibration frequency of 75

Hz.

[104]

Magnetoelectrical Properties Characterization

For the measurements of the magnetoelectric effect, a home-made system was set up based on VSM, as shown in Fig. Thus, the electric field and magnetic field can be applied to the sample simultaneously.

It is necessary to calibrate the measurements to exclude offset the influence of the substrate, electrodes, and wires. Thus we can record the magnetic response of YIG film only is recorded.

FE and AFE Ceramic Substrates

The AFE bulk materials have been developed for decades, but most of them are ceramics. Several researchers have tried to fabricate AFE single crystals, [START_REF] Hlinka | Soft antiferroelectric fluctuations in morphotropic PbZr1-xTixO3 single crystals as evidenced by inelastic x-ray scattering[END_REF][START_REF] Li | Electric field induced metastable ferroelectric phase and its behavior in (Pb, La)(Zr, Sn, Ti)O3 antiferroelectric single crystal near morphotropic phase boundary[END_REF][START_REF] Luo | Progress in lead-based ferroelectric and antiferroelectric single crystals: composition modification, crystal growth and properties[END_REF] but the preparation process of AFE single crystals is not mature, and they have not been widely used in other fields. Therefore, the AFE ceramic is a good choice for us to use in ME heterostructures. As a comparison, the FE ceramic is also selected in this study.

PMN-PZT FE Ceramic Substrates

Results and Discussion

The PMN-PZT ceramics are in the ferroelectric phase with the coexistence of tetragonal and rhombohedra structures. [START_REF] Wang | Effect of PMN content on the phase structure and electrical properties of PMN-PZT ceramics[END_REF] The Curie temperature of PMN-PZT ceramic is 197 ℃. plane strain peak is 0.08%. The large electric-field induced strain change serves to promote the strain-mediated magnetoelectric effect in PMN-PZT-based multiferroic heterostructure.

Summary

The 0.75Pb(Mg1/3Nb2/3)O3-0.25Pb(Zr0.48Ti0.52)O3 ceramic is chosen as the piezoelectric phase to couple with FM films in this work due to its high E-induced strain.

The composition of 0.75PMN-0.25PZT is around the morphotropic phase boundary, where excellent piezoelectric performance exists. The 0.75PMN-0.25PZT ceramics were prepared by solid-state reaction, and average roughness of 17 nm can be reached in polished PMN-PZT ceramic substrates. The EC of PMN-PZT ceramics is as low as 4.2 kV/cm, while the E-induced in-plane and out-of-plane strain can reach 0.08% and -0.27% at EC, respectively.

PLZST AFE Ceramic Substrates

Introduction

The distinctive electromechanical coupling properties of AFE materials have attracted increasing attention ever since the first AFE compound PbZrO3 was discovered by Kittel in 1951. [START_REF] Kittel | Theory of antiferroelectric crystals[END_REF] A typical example is the lightly lanthanum-doped, lead zirconate stannate titanate (PLZST) system. The PLZST system has displayed extremely high electric-field-induced longitudinal and volume strains, especially in the compositions close to the morphotropic phase boundary between the antiferroelectric and ferroelectric phases, as marked in Fig. 4.7. It has been reported that PLZST exhibits large E-field-induced strain (e.g., maximum 0.87% for bulk ceramics and 0.49% for thin films), [START_REF] Pan | Field-forced antiferroelectric-toferroelectric switching in modified lead zirconate titanate stannate ceramics[END_REF][START_REF] Sharifzadeh Mirshekarloo | Large strain and high energy storage density in orthorhombic perovskite (Pb0.97La0.02)(Zr1-x-ySnxTiy)O3 antiferroelectric thin films[END_REF][START_REF] Pan | Large displacement transducers based on electric field forced phase transitions in the tetragonal (Pb 0.97 La 0.02 ) (Ti,Zr,Sn)O 3 family of ceramics[END_REF] potentially enabling a range of device applications, such as actuators, sensors, electrocaloric devices, energy harvesting, and storage devices.

Results and Discussion

The P-E loop and out-of-plane S-E curve of PLZST ceramic given in Fig. 4.10

shows a typical AFE characteristic. The forward and backward switching fields of PLZST are 5 kV/mm and 2.5 kV/mm, respectively. Once the electric-field-induced phase transition between the tetragonal AFE phase and the rhombohedral FE phase happens, a large strain change is generated simultaneously. The maximum strain value along out-of-plane direction can reach 0.2 %, whereas a considerable strain along inplane can also be expected, which provides opportunities for strain-mediated CME coupling effect.

Summary

In this part of the work, the Pb0.97La0.02Zr0.6Sn0.3Ti0.1O3 (PLZST) ceramics, as a typical AFE material in the compositions close to the morphotropic phase boundary between the antiferroelectric and ferroelectric phases, [START_REF] Chen | Preparation of (Pb, La)(Zr, Sn, Ti)O3 antiferroelectric ceramics using colloidal processing and the field induced strain properties[END_REF] were prepared to be the AFE ceramic substrate for its high electric-field-induced strain. The AFE-FE switching field of PLZST is 5 kV/mm with a maximum strain of 0.2 %.

Conclusions

The PMN-PZT FE and PLZST AFE ceramic substrates have been prepared and characterized in this chapter. The ceramic substrates were fabricated by the conventional solid-state reaction method, and the ceramics were machined and wellpolished to reach a low average roughness (< 20 nm) despite the inevitable defects in ceramics. The existence of these defects (holes) may have an influence on the CME coupling performances. The EC of PMN-PZT ceramics is as low as 4.2 kV/cm, while the E-induced in-plane and out-of-plane strain can reach 0.08% and -0.27% at EC, respectively. The AFE-FE switching field of PLZST is 5 kV/mm with a maximum strain of 0.2 %.

The the phase transition for atomic rearrangements and zero polarization and strain in a relaxed state. On the one hand, the significant strain of AFEs, which can reach 0.87 % in some AFE ceramics, make them extremely attractive for strain-mediated ME coupling. [START_REF] Guo | Large Electric Field-Induced Strain and Antiferroelectric Behavior in (1-x )(Na 0.5 Bi 0.5 )TiO 3 -x BaTiO 3 Ceramics[END_REF] On the other hand, the non-remnant state of AFEs results in no prepoling process which shows good compatibility with various applications. [START_REF] Zhou | Antiferroelectric materials, applications and recent progress on multiferroic heterostructures[END_REF] Concerning AFEs, a large strain can be obtained once the electric field is higher than the switching field, which provides a large practical operating voltage range for strains.

So, the distinct characteristics of AFE materials give an additional opportunity to the electric field tuning of magnetism. However, very few works are dedicated to ME coupling of FM/AFE heterostructures, which seems to be interesting and promising in multiferroic composites. [START_REF] Zhou | Strong non-volatile voltage control of magnetism in magnetic/antiferroelectric magnetoelectric heterostructures[END_REF][START_REF] Zhou | Magnetoelectric coupling in antiferroelectric and magnetic laminate composites[END_REF] In order to observe the distinct coupling effect, the AFE PLZST ceramic bulk is selected to couple with FM layers, and the voltage-controlled magnetization switching and magnetic susceptibility variations are studied in different heterostructures.

In strain-mediated CME heterostructures, a strong coupling between elastic and magnetic domains is required for the magnetic component. [START_REF] Wan | Electric-field-induced magnetization in Pb(Zr,Ti)O3/Terfenol-D composite structures[END_REF] The Ni-Mn-Ga (NMG)

alloys have attracted much attention as they combine both magnetic and elastic ordering.

These properties make it possible to manipulate their ferroelastic domains by applying a magnetic field or mechanical stress. [START_REF] Li | Stress-induced change of magnetization in a Ni-Mn-Ga single crystal under magnetomechanical training[END_REF] These characteristics provide opportunities to induce a large magnetization change in NMG by an external mechanical stimulus, which is suitable for strain-mediated CME coupling.

Mirshekarloo and Wang et al. have studied NMG/AFE bilayers in the past, and the strain of austenite-martensite transformation in the NMG top layer can lead to AFE-FE phase transition in AFE films. [START_REF] Mirshekarloo | Ferroelastic strain induced antiferroelectric-ferroelectric phase transformation in multilayer thin film structures[END_REF][START_REF] Zhou | Magnetoelectric coupling in antiferroelectric and magnetic laminate composites[END_REF][START_REF] Wang | Coupling of ferroelastic strain and ferroelectric phase transition in NiMnGa/Pb0.97La0.02(Zr0.95Ti0.05)O3 bilayered films[END_REF] They have demonstrated the strain-induced AFE-FE phase transition, as well as strain-induced polarization change, in NMG/AFE heterostructures. However, the strain-mediated magnetization switching has never been studied in these structures.

Thus, evoked by the previous study of NMG/AFE heterostructure, the E-induced magnetization would be revealed in this part. The NMG film was considered for the FM phase as well as the electrode because of its strong combination of magnetic and elastic orderings and good conductivity. [START_REF] Li | Stress-induced change of magnetization in a Ni-Mn-Ga single crystal under magnetomechanical training[END_REF] 5.2 Preparation of NMG/PLZST/NMG

Film Growth

The NMG films were deposited using a "Plassys" RF sputtering system with a nominal alloy-composition target of Ni0.5Mn0.33Ga0.17. The sputtering condition and annealing process of NMG films were referred to from reported work. [START_REF] Wang | Coupling of ferroelastic strain and ferroelectric phase transition in NiMnGa/Pb0.97La0.02(Zr0.95Ti0.05)O3 bilayered films[END_REF][START_REF] Figueiras | Study of Ni2-Mn-Ga phase formation by magnetron sputtering film deposition at low temperature onto Si substrates and LaNiO3∕Pb(Ti,Zr)O3 buffer[END_REF] The details of the sputtering condition are given in Table 5.1. After the deposition, the samples were annealed in a furnace to be crystallized.

The annealing process is given in Table 5.2. To explore the annealing temperature of NMG films, three NMG/Si samples were prepared and annealed at 300, 400, and 500 ℃, respectively.

ME Heterostructure Fabrication

The NMG films with a thickness of 700 nm (in this case, the ME signal is too weak to detect when the thickness of NMG is low) were deposited on the two sides of PLZST substrates by above mentioned sputtering and annealing process. Since NMG film is conductive, the NMG film is worked as electrodes as well. Thus the NMG/PLZST/NMG multiferroic heterostructure was fabricated, as illustrated in Fig.

5.3.

The NMG polycrystalline film is isotropic, so the direction of the in-plane magnetic field does not matter and can be indifferently along the x or y axis.

Conclusions

In this part of the work, the in-plane CME effect and E-field controlled magnetization based on a composite consisting of NMG films and PLZST AFE ceramic have been investigated. It is remarkable that first, the evolution of (ΔM/M0) as a function of E is in good agreement with the ceramic strain evolution and secondly that the maximum ΔM/M0 variation, which attains a maximum value of 15%, is reached at 0 Oe (without bias magnetic field). The CME coefficient is small, probably due to the defects of NMG film on PLZST ceramic. Because there are inevitable grain gaps in ceramic substrates, many defects were observed in NMG film at the macroscale, and these defects influence the magnetic performance. The study of E-field tuning of magnetism in YIG/AFE heterostructure has not been reported yet, to the best of the authors' knowledge. In this part, the ME composites of YIG/PLZST has been fabricated, and the in-plane and the out-of-plane voltage induced magnetism has been carried out. The YIG/PMN-PZT heterostructure has also been fabricated, and the E-field controlled magnetization switching and E-tunable susceptibility would be reported.

Preparation of YIG/AFE and YIG/FE Heterostructures

YIG film Growth

The YIG films were prepared by using a "Plassys" RF sputtering system with a YIG powder target. The sputtering condition and annealing process of YIG films have been explored by several researchers, [START_REF] Lian | Annealing effects on the microstructure and magnetic properties of Y3Fe5O12 films deposited on Si/SiO2 substrates by RF magnetron sputtering[END_REF][START_REF] Kang | Magnetic properties of YIG (Y3Fe5O12) thin films prepared by the post annealing of amorphous films deposited by rf-magnetron sputtering[END_REF][START_REF] Stognij | Growth and spin-wave properties of thin Y3Fe5O12 films on Si substrates[END_REF] and in particular in the group. They have demonstrated that the post-annealing process has strong effects on the quality of YIG films. The details of the sputtering condition are given in Table 6.1. After the deposition, the samples were annealed in a furnace to be crystallized.

The reported crystallization temperature of the YIG film is 700~750 ℃. [START_REF] Lian | Annealing effects on the microstructure and magnetic properties of Y3Fe5O12 films deposited on Si/SiO2 substrates by RF magnetron sputtering[END_REF][START_REF] Kang | Magnetic properties of YIG (Y3Fe5O12) thin films prepared by the post annealing of amorphous films deposited by rf-magnetron sputtering[END_REF] The annealing process is given in Table 6.2. The YIG films were annealed at 700 and 750 ℃. Cooling Rate (℃/min) 1

Atmosphere Air

The YIG films were firstly deposited on SiO2/Si substrates to check the sputtering condition and annealing process. Figure 6.1 shows the XRD patterns and magnetic hysteresis loops of YIG films on SiO2/Si substrates with the different annealing process.

From Fig. 6.1 (a), it can be observed that the YIG films cannot be crystallized on SiO2/Si substrates when the annealing temperature is 700 ℃, but the crystallization can be achieved at 750 ℃. In Fig. 6.1 (b), the well-crystallized YIG film shows good magnetic properties. The high saturated magnetization (MS =105 emu/cm 3 ) and low magnetic coercive field (HC =43 Oe) can be reached in YIG films, which reveals the feasibility of the preparation process of YIG film. In order to apply the electric field, a soldering process was involved in connecting conductive wires with Pt electrodes. Thus the magnetic field and electric field can be applied to the sample simultaneously. The YIG film being polycrystalline it demonstrates isotropic in-plane magnetic properties, thus the in-plane magnetic field direction is parallel to the substrate plane.

The in-plane M-H loop of YIG/Pt/PLZST/Pt heterostructure is shown in Fig. 6.5.

The YIG film presents good magnetic properties on PLZST ceramic substrate, the same as YIG on PMN-PZT ceramic substrate. The values of the saturate magnetization (MS), 60 remnant magnetization (Mr), coercive magnetic field (HC) and magnetic susceptibility (χ) in YIG/Pt/PLZST/Pt are given in Table 6.3. The M-H loops of YIG/Pt/PLZST/Pt heterostructure while an electric field is applied across the PLZST ceramic are measured in order to investigate the magnetoelectric coupling in a static electric state. The M-H loops with different electric fields and HC-E curve are given in Fig. 6.6 (a) and (b). When the external electric field is lower than AFE-FE switching field of PLZST, the M-H loops coincide with the one without an electric field. Once the electric field is higher than the AFE-FE switching field, the remnant magnetization decreases sharply, and the M-H loops become difficult to be saturated. In this case, an induced magnetic coercive field shift is generated in response to the application of an electric field. The strain driven by the electric-fieldinduced AFE-FE phase transition is transferred through Pt electrode to YIG film, which leads to magnetic anisotropy transformation in the YIG film. Additionally, the HC-E and χ-E curves clearly display the familiar horn-like shape but in the opposite direction, which is attributed to the correlation between strain and electric field. The shape of HC-E and χ-E curves demonstrate the effective control of magnetocrystalline anisotropy transformation by E-field via strain mediated mechanism.

In order to investigate the magnetoelectric coupling in a static magnetic state, the M-E curves were measured while applying an electric field across the PLZST ceramic.

Before the measurements, the influences of the ceramic and the electrodes were is, of course, in relation to bias H applied to the YIG film during the measurement. For example, in Fig. 6.8 (H = 0 Oe), M0i associated is 37.8 emu/cm 3 , while M0i equals to 45.2 emu/cm 3 in Fig. 6.8 (H = 30 Oe). This difference of (M0i-M0f) is due to the irreversible demagnetization effect induced by the strain. The strain can easily destroy the alignment of well-ordered magnetic domains and favors to rotate magnetic moments toward the opposite direction. However, when the strain is removed, not all the magnetic domains can recover to its initial state. The unrecoverable part leads to the difference of (M0i-M0f). Thus the difference of (M0i-M0f) only exists in the first half electrical cycle of measurement.

As the bias magnetic field decreases, the ratio of the strain-controlled magnetization change (∆M/M0f) as a function of the bias magnetic field is shown in Fig. 6.9 (a). It can be seen that ∆M/M0f gradually decreases with the increase of the applied bias magnetic field, and there is a maximum value of 17.5% at the zero magnetic fields.

This phenomenon indicates that the influence of electric field induced strain on YIG film is maximized when the bias magnetic field is 0 Oe, which is of great significance for the development of future magnetoelectric devices. In contrast, in the high magnetic field (30 Oe), the mechanical deformation of ceramics has a relatively limited regulation of the magnetization of YIG films, with only 11% of the magnetization changes. The ratio (M0i-M0f)/ M0f depends on the magnetic state of the YIG film. This ratio, as shown in Fig. 6.9 (b), gets its maximum at 0 Oe and then decreases quickly, as a function of H. The maximums of αCME for three different bias magnetic fields (0, 15, 30 Oe) are reported in Fig. 6.12. The αCME decreases with a bias magnetic field increasing, and the largest value of αCME is obtained with a bias magnetic field of 0 Oe. The fact that the best performance of CME coupling is obtained without the magnetic field is a really important asset in terms of design when targeting ME devices conception. CME coefficient is obtained without the application of bias magnetic field, which is a strong asset in terms of cost and ease of procedures. The high relative susceptibility change reveals a potential opportunity of AFE materials to develop future ME tunable devices.

Out-of-plane CME Effect in YIG/Pt/PLZST/Pt Heterostructure

In order to fully understand the CME Effect in YIG/Pt/PLZST/Pt heterostructure, the out-of-plane E-field induced magnetism was also carried out. For these measurements, the magnetic field is vertically applied to the YIG/Pt/PLZST/Pt heterostructure. The electric-field dependence of out-of-plane M-H loops of YIG/Pt/PLZST/Pt heterostructure are illustrated in Fig. 6.13. It can also be observed that the out-of-plane saturate magnetic field of YIG film is much higher than in-plane results. The YIG polycrystalline film has an "easy plane" on the PLZST substrate plane that more energy is needed for YIG film to be saturated along out-of-plane direction. Under an application of electric field, the M-H loops weakly varies with the electric-field-induced strain.

The out-of-plane magnetization change as a function of the electric field at 500 and 1000 Oe in YIG/Pt/PLZST/Pt heterostructure are given in Fig. 6.14. The horn-like shape of M-E curves can be clearly observed, even the strain-induced magnetic variations are small. The out-of-plane strain favors the magnetization process. With the bias magnetic field increasing, the strain-induced magnetization effect is also hindered by the bias magnetic field. The in-plane and out-of-plane M-E curves of YIG/Pt/PLZST/Pt heterostructure are given in Fig. 6.15. The out-of-plane strain favors the magnetization process, while inplane strain leads to the demagnetization process. The electric field induced in-plane magnetization variations are much larger than out-of-plane ones. These differences can also be explained by the strain-induced magnetic domain motion is hindered in thickness-direction due to its low dimension. However, large strain-induced magnetic variation can be achieved because of its large dimension in width-and length-direction.

Summary

In summary, the E-field controlled magnetism in an AFE-based multiferroic heterostructure of YIG/PLZST is reported. A comprehensive study of in-plane magnetization, coercive magnetic field, magnetic susceptibility, and CME coefficient response to an electric field is achieved in this part. The electric field induced phase transition strain in AFE substrate plays a crucial role in ME coupling, and its unique zero-remnant state gives us extra advantages to avoid the pre-poling process. The maximum in-plane αCME is 11.6 × 10 -8 s/m without a magnetic field, and the in-plane ∆χ/χ0 can reach 33% with a low magnetic field of 10 Oe. The large susceptibility variation possibility given by this composite, YIG/Pt/PLZST/Pt, shows great potential for ME tunable devices. This E-field control of magnetism is originated from the Efield induced magnetic anisotropy change. The tensile out-of-plane strain generated in PLZST ceramic leads the magnetic anisotropy of YIG film to tilt to out-of-plane direction, thus results in in-plane demagnetization and out-of-plane magnetization effect with an application of the electric field.

YIG/Pt/PMN-PZT/Pt heterostructure

The YIG/FE heterostructures have been intensively studied in the past years, and various systems have been developed. In this part, the CME coupling effect of YIG/Pt/PMN-PZT/Pt was studied as a primary objective to compare the FE-based ME coupling effect with Chapter 6.2. The in-plane and out-of-plane CME coupling effect would be discussed in YIG/Pt/PMN-PZT/Pt heterostructure.

In-plane CME Effect in YIG/Pt/PMN-PZT/Pt heterostructure

For these results, the applied magnetic field is parallel to the YIG/Pt/PMN-PZT/Pt plane. Figure 6.17 shows the M-H loops of YIG/Pt/PMN-PZT/Pt heterostructure. As mentioned before, the YIG film on PMN-PZT ceramic is polycrystalline, indicating that the YIG film is isotropic in the plane. Good magnetic properties of YIG film can be obtained on PMN-PZT ceramic substrate, where the MS, Mr, HC, and χ are given in Table 6.5. The YIG films on PLZST and PMN-PZT are prepared by the same process, but the values in Table 6.5 are not close to the ones in Table 6.3. It can be explained by the differences of the ceramic substrates. Since the ceramic is a complex polycrystalline system and has many intrinsic holes, every piece of ceramic can be different. Even all the ceramic substrates are polished into an average roughness of several nanometers, the big grain gaps also exist. Also, they would affect the magnetic properties of YIG films and maybe the CME coefficient. As mentioned in Chapter 6.2.1, the YIG/Pt/PLZST/Pt heterostructure was fully magnetized at the same high bias magnetic field (2 T) to ensure the same initial state.

Since PMN-PZT is a ferroelectric material, it presents remnant state after an application of the electric field. To eliminate the influence of electrical history, the application of an electric field cycle starts from a value of +12 kV/cm to conform to the strain and polarization states of PMN-PZT. The change of the electric field is as follows: +12 → -12 → +12 → 0 kV/cm. Then the second cycle of the electric field can be applied as: 0 s/m in the absence of a bias magnetic field at a low electric field of 4.2 kV/cm. Thus these results make it possible to achieve a strong ME effect under a small electric field and zero magnetic field that promotes the realization of power-efficient ME devices. where the magnetization can be easily switched. Thus the αCME reaches the peak at this point, which is probably attributed to the easy magnetization switching at HC. When the bias magnetic field is higher than HC, the electric-field-induced demagnetization process is hindered by a large bias magnetic field that results in a low CME coefficient.

→ -12 → +12 → 0 kV/cm.
Since the E-induced strain plays a critical role in the ME coupling effect, the outof-plane S-E curves of PMN-PZT ceramic measured by cycling electric field with different amplitudes are depicted in Fig. 6.23. The in-plane strain should show a similar evolution as the out-of-plane strain. When the amplitude of E is much higher than EC, a whole butterfly-like curve appeared because of complete ferroelectric domain switching. While E ≤ EC, the shuttle-like curve was observed due to incomplete ferroelectric domain switching, which results in large remnant strain. Thus, two different strain states are generated when E ≤ EC, leading to two magnetization states in YIG film. The M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure are asymmetric, and there is a slope. To remove the influence of electric load history, the YIG/Pt/PMN-PZT/Pt heterostructure should be poled before the measurements. If it was poled at an electric field of +12 kV/cm for 3min, then the M-E curves would incline to E+, as shown in Fig. 6.25 (a). In Fig. 6.25 (b), the M-E curves would incline to E-when the YIG/Pt/PMN-PZT/Pt heterostructure was poled at an electric field of -12 kV/cm first. Thus, the inclined direction of asymmetric M-E curves is depended on the electric field history before the measurements. show the great potential in the application of energy-efficient ME random access memory devices. While the amplitude of E equals 4 kV/cm, the slim shuttle-like M-E curve indicates large and linear E-tunable susceptibility for RF/microwave magnetic devices.

In order to simulate the real memory write process, a dynamic voltage-pulseinduced magnetization switching in YIG/PMN-PZT, is shown in Fig. 6.27. With the zero magnetic bias field, the magnetization response to the electric field shows a reversible hysteretic type behavior, as illustrated in Fig. 6.27 (a). The remnant magnetization was reversibly switched from state "1" to state "0" with a magnitude of magnetization ratio of 6.2%. Furthermore, memory-type dynamically switching of magnetization between the two magnetization states is demonstrated in Fig. 6.27 (b).

The electric field pulses of 4.6 kV/cm and -4.6 kV/cm are applied alternately at a period of 60 s. As an E-field pulse of 4.6 kV/cm has been applied, magnetization in YIG rises and remains at a magnetization ratio of 54.4%, which means a switch from the state "1" to state "0". Similarly, the magnetization was reduced and stayed at a magnetization ratio of 48.2%, as an E-field pulse of -4.6 kV/cm has been applied, indicating the switching from "0" to "1."Therefore, dynamic E-field-pulse-induced magnetization switching in YIG was realized with zero magnetic field and a low electric field of 4.6 kV/cm. This memory-type of magnetization switching would result in the non-volatile magnetization tuning with excellent energy efficiency. χ-E curve is observed, and the χ varies in the range of 15.5 and 32.9. The large E-tunable susceptibility ratio can reach ∆χ/χ0+ =55.5% at 20 Oe and with a cycling electric field of 4 kV/cm. Other than conventional magnetic components or devices, E-tunable devices with larger tunability and lower energy consumption, where tunability is achieved by adjusting the electric field, can be developed. [START_REF] Liu | Voltage tuning of ferromagnetic resonance with bistable magnetization switching in energyefficient magnetoelectric composites[END_REF] Thus two interesting multiferroic applications can be developed depending on YIG/Pt/PMN-PZT/Pt heterostructures. On the one hand, the voltage-pulse reversible magnetization switching of YIG/Pt/PMN-PZT/Pt heterostructures is realized at 0 Oe with an electric field impulse of ±4.6 kV/cm, which shows the feasibility of powerefficient E-field controlled nonvolatile memories. On the other hand, the linear Etunable susceptibility change can be realized at 20 Oe and 4 kV/cm, and the tunable susceptibility can reach ∆χ/χ0+ =55.5%, indicating the potential of power-efficient magnetoelectric tunable devices.

Out-of-plane E-field controlled Magnetism in YIG/Pt/PMN-PZT/Pt heterostructure

For these measurements, the magnetic field is vertically applied to the YIG/Pt/PMN-PZT/Pt heterostructure.

The electric-field dependence of out-of-plane M-H loops of YIG/Pt/PMN-PZT/Pt heterostructure are illustrated in Fig. 6.29. With the absence of an electric field, the outof-plane M-H loop shows different shapes from the in-plane M-H loop. It can be observed that the out-of-plane saturate magnetic field of YIG film is much higher than in-plane results, indicating that the YIG film has an "easy plane" in the PMN-PZT substrate plane that more energy is needed for YIG film to be saturated along out-ofplane direction. Under an application of electric field, the M-H loops weakly vary with the electric-field-induced strain. The minor electric field induced magnetic variations demonstrate that the strain-induced magnetization or demagnetization effect is weak along the out-of-plane direction. These results can be explained by the strain-induced magnetic domain motion is hindered in thickness-direction due to its low dimension.

However, large strain-induced magnetic variation can be achieved because of its large dimension in width-and length-direction. The out-of-plane magnetization change as a function of the electric field at 500 and 1000 Oe in YIG/Pt/PMN-PZT/Pt heterostructure are given in Fig. 6.30. The butterfly-like shape of M-E curves can be observed; even the strain-induced magnetic variations are small. The compressive out-of-plane strain favors the magnetization process, while tensile out-of-plane strain leads to the demagnetization process, which agrees with in-plane results. With the bias magnetic field increasing, the strain-induced magnetization or demagnetization process is also hindered by the bias magnetic field. The typical in-plane and out-of-plane M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure are given in Fig. 6.31. Accounting to the in-plane and out-of-plane S-E curves of PMN-PZT, given in Fig. 4.6, it can be concluded that the compressive strain favors the magnetization process, while tensile strain leads to the demagnetization process in the YIG film. The electric field induced in-plane magnetization variations are much larger than out-of-plane ones. These differences can be explained by the strain-induced magnetic domain motion is hindered in thickness-direction due to its low dimension. However, large strain-induced magnetic variation can be achieved because of its large dimension in width-and length-direction. When the electric field induces an in-plane compressive strain and an out-of-plane tensile strain at the same time, the anisotropy axis will tilt to in-plane direction. Thus the in-plane magnetization (Mip) increases, and out-of-plane magnetization (Moop) decreases. In contrast, when the electric field induces an in-plane tensile strain and an out-of-plane compressive strain, the anisotropy axis would tilt to the out-of-plane direction. Thus the in-plane magnetization (Mip) reduces, and out-of-plane magnetization (Moop) rises.

Summary

In this part, we have investigated the in-plane E-field controlled magnetism in YIG/Pt/PMN-PZT/Pt heterostructure. The electric-field-dependent M-H loops and χ-H curves reveal the strain-mediated transformation of magnetocrystalline anisotropy in the YIG film. As a result, the E-field controlled Mr, HC, and χ change correspondingly with the PMN-PZT ceramic's strain evolution. The CME coefficient can reach 16.95 × 10 -8 s/m in the absence of bias magnetic field at EC of PMN-PZT ceramic, and the maximum of CME coefficient (18.15 × 10 -8 s/m) is observed under the magnetic field of 25 Oe, demonstrating strong converse magnetoelectric coupling effect in 84 YIG/Pt/PMN-PZT/Pt heterostructure. When the application of electric field amplitude is close EC of PMN-PZT, the E-induced shuttle-like strain evolution of PMN-PZT favors two possible applications in YIG/Pt/PMN-PZT/Pt heterostructure. On the one hand, two distinct remnant strain states lead to two stable magnetization states of the YIG film. The voltage-impulse reversible magnetization switching of YIG/Pt/PMN-PZT/Pt heterostructures is realized at 0 Oe with an electric field impulse of ±4.6 kV/cm, which shows the feasibility of power-efficient E-field controlled nonvolatile memories.

On the other hand, the linear E-tunable susceptibility change can be realized at 20 Oe and 4 kV/cm, and the tunable susceptibility can reach ∆χ/χ0+ =55.5%, indicating the potential of power-efficient magnetoelectric tunable devices.

The E-field control of magnetism in YIG/Pt/PMN-PZT/Pt heterostructures is achieved by E-field induced magnetic anisotropy change. When the electric field induces an in-plane compressive strain and an out-of-plane tensile strain in PMN-PZT ceramic at the same time, the anisotropy axis of YIG film will tilt to in-plane direction.

In contrast, when the electric field induces an in-plane tensile strain and an out-of-plane compressive strain in PMN-PZT ceramic, the anisotropy axis of YIG film would tilt to the out-of-plane direction.

Conclusions: comparison of FE (AFE)/YIG performances

The in-plane and out-of-plane E-field control of magnetism in YIG/AFE and YIG/FE (YIG/Pt/PLZST/Pt and YIG/Pt/PMN-PZT/Pt) heterostructures are studied in this part of work. The CME coupling effect in YIG/AFE and YIG/FE does share the same strain-mediated coupling mechanism, as given in following:

1. The E-field controlled Mr, HC, and χ of YIG films change correspondingly with the substrate's strain evolution (S-E curves), indicating the effectiveness of strainmediated CME coupling effect.

2. The strain-induced magnetic anisotropy transformation of YIG film results in the magnetization or demagnetization effect in the plane or along the OOP direction with an application of the electric field.

However, there are some differences in performances due to the different strain evolutions of the substrates under the electric field. The AFE materials display zero 85 remnant states; thus, in AFE-based ME composites, the pre-poling process is not necessary. The AFE-FE switching field of PLZST is 50 kV/cm with a maximum strain of 0.2 %. The FE materials show hysteresis loops, and the electric load history influences the following results. The EC of PMN-PZT ceramics is as low as 4.2 kV/cm, while the E-induced in-plane and out-of-plane strain can reach 0.08% and -0.27% at EC, respectively. When the electric field is lower than EC, the S-E curves are shuttle-like, and there are two remnant strain states.

1. The large E-induced phase transition strain in AFE substrate leads to maximum in-plane αCME (11.6 × 10 -8 s/m) at 0 Oe in YIG/Pt/PLZST/Pt heterostructure under an electric field of 50 kV/cm. The in-plane CME coefficient in YIG/Pt/PMN-PZT/Pt heterostructure can reach 16.95 × 10 -8 s/m in the absence of bias magnetic field at EC of PMN-PZT ceramic, and the maximum of in-plane CME coefficient (18.15 × 10 -8 s/m) is observed under the magnetic field of 25 Oe (@HC). However, it is necessary to point out that even though the maximum of αCME is obtained at a low electric field of 4.2 kV/cm, but a pre-polling process to PMN-PZT substrate is essential to obtain this high αCME. Otherwise, a large in-plane strain of -0.27% cannot be obtained at EC. Good CME coefficients of YIG/Pt/PLZST/Pt and YIG/Pt/PMN-PZT/Pt show that AFE and FE materials are both potential candidates in ME composites. However, in general, the switching field of AFE materials is always very high, which may be not suitable to develop AFE-bulk ME composites. The required high voltage of AFE bulks would greatly restrict their applications in ME composites. The reduction of the thickness in AFE materials should be a solution to avoid high voltages. Thus, AFEbased ME thin films are a possible way to utilize the advantages (no pre-poling process) of AFE materials and avoid high voltages, which is interesting for volatile memory devices.

2. When the application of electric field amplitude is close to EC of PMN-PZT, the voltage-impulse reversible magnetization switching and the linear E-tunable susceptibility change can be realized in YIG/Pt/PMN-PZT/Pt heterostructures because of the shuttle-like S-E curves of PMN-PZT substrate. These features of FE materials give us another way to develop FE-based ME composites with low power consumption.

The reversible voltage-impulse reversible magnetization switching is realized at 0 Oe and with the application of electric field pulse of which the amplitude is close to EC of PMN-PZT, implying great potential in non-volatile memory devices.

3. The maximum of in-plane ∆χ/χ0 in YIG/Pt/PLZST/Pt heterostructure is 33% at 10 Oe and with an electric field cycle from 0 to 50 kV/cm, and a linear E-tunable susceptibility change is realized. In YIG/Pt/PMN-PZT/Pt heterostructures, the linear Etunable susceptibility change can be realized at 20 Oe and with an electric field cycle from -4 to +4 kV/cm, and the tunable susceptibility can reach ∆χ/χ0+ =55.5%. Once again, the AFE-and FE-based ME composites show good potential in E-tunable microwave devices. The difference in the values in ∆χ/χ0 may be due to the differences in strain values. Again, the reduction of the thickness in AFE materials should be a solution to avoid high voltages.

The E-field controlled magnetism in YIG/AFE and YIG/FE demonstrated that the electric-field induced strain is responsible for the plays a vital role in the CME coupling effect. The difference in the relationship of strain and electric field in ceramic substrates favors the difference in the CME coupling effect of YIG/AFE and YIG/FE. The interfaces quality plays, of course, a key role in the coupling efficiency, we do not take into account this parameter in our studies.

CME Coupling Effect of [(TbCo2)/(FeCo)]20 Films on AFE and FE Ceramics

Various materials have been used as the ferromagnetic/ferrimagnetic phases to couple with FE phases. In the last chapter, the CME coupling effect between the ferrimagnetic phase (YIG) and AFE or FE phases has already been demonstrated. In general, ferromagnetic materials, especially the 3d-transition metals (Fe, Co, and Ni) and alloys (permendur, Terfenol-D, Galfenol), present a more excellent magnetic performance at low frequency. They are also essential parts of designing FM/FE composites, in which the excellent ME coupling effect has already been demonstrated.

The rare-earth metal Tb modified Co/Fe alloy shows excellent magnetic properties, high magnetostriction, and large saturated magnetization when compared with ferrite YIG, [START_REF] Dusch | Stress-mediated magnetoelectric memory effect with uni-axial TbCo2/FeCo multilayer on 011-cut PMN-PT ferroelectric relaxor[END_REF][START_REF] Hall | Magnetic anisotropy and magnetostriction of ordered and disordered cobalt-iron alloys[END_REF][START_REF] Jung | Soft anisotropic high magnetization Cu/FeCo films[END_REF] which makes it an excellent ferromagnetic material in ME composites. In this chapter, the ferromagnetic materials (Tb modified Co/Fe alloy) are used to couple with AFE and FE phases to understand the AFE-and FE-based CME coupling effects. Amorphous [(TbCo2)/(FeCo)]20 (TCFC) film was chosen to be another FM layer due to its high perpendicular magnetic anisotropy, which allows us to study the CME effect along different in-plane axes. [START_REF] Rahman | Compositional dependence of magnetoresistance in TbFeCo amorphous film[END_REF][START_REF] Tiercelin | Magnetoelectric effect near spin reorientation transition in giant magnetostrictive-aluminum nitride thin film structure[END_REF] The previous work of [(TbCo2)/(FeCo)]x (TCFC) films on PMN-PT single-crystal slabs has shown great potential in voltagecontrolled magnetoelectric read and write operations when focusing at random access memory devices. [START_REF] Dusch | Stress-mediated magnetoelectric memory effect with uni-axial TbCo2/FeCo multilayer on 011-cut PMN-PT ferroelectric relaxor[END_REF][START_REF] Tiercelin | Enhanced magnetoelectric effect in nanostructured magnetostrictive thin film resonant actuator with field induced spin reorientation transition[END_REF][START_REF] Klimov | Magnetoelectric write and read operations in a stress-mediated multiferroic memory cell[END_REF] TCFC films show many advantages in ME composites, but the study of E-field tuning of magnetism in TCFC/AFE heterostructures has not been reported yet. In this chapter, the TCFC/PLZST heterostructure was carried out, and the CME coupling effect has been studied. Since the TCFC is anisotropic in the plane, it is interesting to explore the differences of the CME coupling effect along different inplane axes (easy axis and hard axis). To further understand the CME coupling effect of TCFC films on ceramic substrates, the TCFC/PMN-PZT heterostructure has also been carried out in this part, and the voltage induced magnetism along different axes has been studied.

Preparation of TCFC/AFE and TCFC/FE Heterostructures

The [(TbCo2)/(FeCo)]20 multilayers were deposited by sputtering on one side of the PLZST and PMN-PZT ceramic substrates layer by layer, where the thickness of each layer is 5 nm, as given in Fig. 7.1 (a). An easy magnetic axis is obtained in the magnetic layer by applying a magnetic field during deposition. This layer also serves as the top electrode for its good conductivity. The final thickness of the multilayers is 200 nm. The dimensions of TCFC film was restricted to be 6 × 6 mm 2 by masks. Since the TCFC films are sensitive to temperature and can lose its magnetic properties at ~100 ℃. A soldering process for wire attachments is not suitable in this case. Then the Au electrodes were deposited on both sides on the heterostructure. The choice of Au electrodes is for the following nano Au wire growth method to attach wires. The easy axis and hard axis of TCFC film are perpendicular, and they are controlled to be in the plane through the deposition process, as shown in Fig. 7.1 (b). The TCFC film is sensitive to temperature that the hot soldering process should avoid being directly on TCFC film. A designed sampler holder is used for wire connection. Figure 7.2 illustrates that the TCFC/PLZST (PMN-PZT)/Au heterostructure is able to connect with wires through a connection of sample holder.

Firstly, the front and back Au electrodes of TCFC/PLZST (PMN-PZT)/Au heterostructure connect to the front and back sides of the sample holder using the nano Au wire growth method, respectively. Then the wires can be attached to the conductive front and back sides of the sample holder through the soldering process.

[(TbCo2)/(FeCo)]20/PLZST/Au Heterostructure

Magnetocrystalline anisotropy is the energy necessary to deflect the magnetic moment in a single crystal from the easy to the hard direction, and it is an intrinsic property of magnetic materials. It can be easily seen by measuring magnetization curves along with different crystal directions. Depending on the crystallographic orientation of TCFC film in the magnetic field, the magnetization reaches saturation in different fields. The magnetization in the easy axis is mostly easy to reach the saturation, while the magnetization in hard-axis is usually saturated in a higher magnetic field. The M-H loops of TCFC/PLZST/Au heterostructure along different directions are given in Fig. 7.3. The TCFC multilayer presents excellent magnetic performance on the ceramic substrate. A saturated magnetization of ~700 emu/cm 3 was observed in the TCFC multilayer, demonstrating its strong ferromagnetism. The TCFC multilayer is a uniaxial film in the plane; thus, different M-H loops were obtained along with different directions. When the magnetic field is along the OOP direction, the TCFC film cannot be saturated with a magnetic field of 2.1 Tesla, indicating there is a strong in-plane anisotropy field that larger energy is needed to deflect the magnetic moment from inplane to OOP direction. The saturated magnetization of TCFC can reach 706 emu/cm 3 , and the HC is 190 Oe. It can be observed that the TCFC film gets saturated with a magnetic field of ~1500 Oe. In TCFC film, the magnetic moments are spontaneously aligned to the easy axis.

CME Effect along Hard Axis

When the external magnetic field is perpendicular to the easy axis, magnetic moments become to rotate. If the amplitude of the external magnetic field is lower than the saturation field of TCFC, parts of magnetic moments are spontaneously aligned to the easy axis, parts are aligned to the hard axis, and the rest parts are in between. In this case, if the electric field induced strain is applied to TCFC film, the orientation of magnetic moments can be easily changed.

The M-E curves of TCFC/PLZST/Au heterostructure along the hard axis are given in Fig. 7.5. Under a specific magnetic field, the sample was also saturated in a high magnetic field of 2.1 Tesla before the measurements to eliminate the historical effect, then the first cycle and second half cycle of electric-field-induced magnetization were recorded. For the first cycle of the M-E loop, the magnetization variations can be divided into two parts, that is the reversible part and the irreversible part. As it has been discussed in Chapter 6.2.1, the irreversible part only occurs in the first M-E loops, and it can be characterized by the difference of the magnetization between the initial state (M0i), at the beginning of the measurements, and the magnetization state at the end of first electrical cycle (M0f) in the first cycle of M-E loop. For the second half of the M-E loop, the magnetization of the initial state and the end state are almost the same. Thus it is reversible.

In Fig. 7.5, when the external magnetic field is lower than HC, the M-E curves show typical horn-like loops, which is in the same shape of the S-E curve of the PLZST the alignment of magnetic moments to the hard axis and leads to a demagnetization process, and the removal of strain (FE-AFE) can lead to a magnetization process. When the external magnetic field is equal to or greater than HC, the application of strain gradually becomes a role in leading to a magnetization process, and the removal of strain leads to a demagnetization process. In the first cycle of the M-E curve at 500 Oe, the magnetization of TCFC first decreases greatly at the AFE-FE switching field, which can be attributed to that the strain destroys the remnant alignment of magnetic moments to the hard axis. When the external magnetic field is 1000 Oe, the applied electric field induced strain of PLZST can rise the magnetization of TCFC, indicating that the strain helps the magnetic moments to align to the hard axis. And the remove of strain can destroy the alignment of magnetic moments and lead to a demagnetization process. In a word, with the application of different amplitudes of the external magnetic field, the applied electric field induced strain can result in different effect, i.e., a demagnetization effect when H < HC, a magnetization effect when H ≥ 1000 Oe and a mixing effect at HC ≤ H < 1000 Oe. A detailed discussion and analysis of this phenomenon will be given in Chapter 7. 8% with an absence of a magnetic field, then it goes up to 1.9% at 500 Oe and reaches the maximum of 2.3% at 1000 Oe. Without the external magnetic field, the αCME is at the negative maximum of -10.3×10 -8 s/m, and the positive maximum of αCME is at 500 Oe with a value of 12.7×10 -8 s/m. in-plane anisotropy field that larger energy is needed to deflect the magnetic moment from in-plane to OOP direction. In this case, the energy of E-induced strain in PLZST maybe not high enough to deflect the magnetic moment from in-plane to OOP direction.

CME Effect along Easy Axis

CME Effect along OOP Direction

Thus the CME coupling effect is too small to be detectable along the OOP direction.

Summary

In this part of the work, the [(TbCo2)/(FeCo)]20/PLZST/Au heterostructure has been prepared, and the E-induced magnetism along different axes has been studied.

Along the easy axis of TCFC film, the E-induced magnetization variation is suppressed by the highly intrinsic crystallographic orientation with the magnetic field. Along the hard axis, the E-field controlled magnetization under different bias magnetic fields has been demonstrated, as well as the CME coefficient. A reverse of M-E curves occurred with the bias magnetic field increasing along the hard axis. The maximum of ∆M/M0 can reach -9.8% with an absence of the magnetic field, and the αCME is at the negative maximum of -10.3×10 -8 s/m at the same time. The positive maximum of αCME is at 500

Oe with a value of 12.7×10 -8 s/m. Along OOP direction, the E-induced magnetism is not visible, which can be attributed to that the energy of E-induced strain in PLZST is not enough to deflect the magnetic moment from in-plane to OOP direction.

[(TbCo2)/(FeCo)]20/PMN-PT/Au heterostructure

The M-H loops of TCFC/PMN-PZT/Au heterostructure along different axes are given in Fig. 7.10. The TCFC multilayer presents excellent magnetic performance on the PMN-PZT substrate. A saturated magnetization of ~600 emu/cm 3 was observed. When the magnetic field is along the OOP direction, the TCFC film cannot be saturated with a magnetic field of 2.1 Tesla, indicating there is a strong in-plane anisotropy field that larger energy is needed to deflect the magnetic moment from in-plane to OOP direction. along the hard axis. A higher magnetic field is needed to deflect magnetic moments from the easy to the hard direction in TCFC film. Thus when the applied magnetic field is low, some parts of the magnetic moments in TCFC film are aligned along the hard axis, and some are not. Then the large strain generated in PMN-PZT ceramic can easily rotate some intermediate magnetic moments and lead to large magnetization variations.

CME Effect along Hard Axis

As given in Fig. 7.11 inset, when the magnetization is not saturated in TCFC film, the distinguishable strain-induced magnetization variations are observed. The M-E curves of TCFC/PMN-PZT/Au heterostructure along the hard axis are carried out under different magnetic fields. Under a specific magnetic field, the sample was saturated in a high magnetic field of 2.1 Tesla before the measurements to eliminate the historical effect, then the first cycle of electric-field-induced magnetization was recorded. Also, several second cycles of electric-field-induced magnetization were also recorded. Thus the M-E curves under a specific magnetic field were finished, then repeated the steps. Figure 7.12 shows how an electric cycle goes. Since the FE ceramic has remnant state after an application of electric field, the applied electric field started from -12 kV/cm to assure that the ceramic was in the same state before the measurements.

- M-E curve at 0 Oe. At 800 Oe, the E-field controlled magnetization variations are tiny, and it is difficult to find a relationship between the E-field induced strain and magnetization. When the external magnetic field keeps increasing, the M-E curve is completely inversed at 1000 Oe, and there are two magnetization peaks around EC.

To further explain the bizarre behavior of M-E curves, the ideal magnetic hysteresis loop along the hard axis is given in Fig. 7.14. For a perfect single crystal, when the applied magnetic field is perpendicular to the easy axis of the material, it presents an oblique-line-like M-H loop, and no remnant magnetization and zero HC.

The single crystal saturates at the anisotropy field (Ha), where Ha=2K/MS, MS is the saturated magnetization, K is the anisotropy coefficient.

If the magnetoelastic coupling constant of a single crystal is negative, and the anisotropy field is H0. According to Equation 2.4, the external tensile strain would lead more magnetic moments to align perpendicularly to the strain direction, and the magnetic moments tend to align along compressive strain direction. In this case, when an E-field induced tensile strain is given, the anisotropy field decreases to H1. While an E-field induced compressive strain is given, the anisotropy field rises to H2. In TCFC/PMN-PZT/Au heterostructure, the magnetoelastic coupling constant of TCFC is negative (-7 MPa). [START_REF] Tiercelin | Enhanced magnetoelectric effect in nanostructured magnetostrictive thin film resonant actuator with field induced spin reorientation transition[END_REF] When an E-field induced tensile strain is applied to the film, the anisotropy field of TCFC should decrease, and the anisotropy field increases when an E-field induced compressive strain is applied. Figure 7.15 illustrates the details of the real M-H loops ng hard axis in TCFC/PMN-PZT/Au heterostructure with an application of the electric field. When an electric field of 4 kV/cm is applied to PMN-PZT ceramic, a large tensile in-plane strain and compressive OOP strain would be generated in the ceramic substrate and be transferred to TCFC film. Figure 7.15 (a) shows how the M-H loop changes with the Efield induced tensile in-plane strain and compressive OOP strain. The slope of the M-H loop varies synchronously with the ideal model, demonstrating the reduction of anisotropy field. When an electric field of -12 kV/cm is applied to PMN-PZT ceramic, a large compressive in-plane strain and tensile OOP strain would be generated in the ceramic substrate and be transferred to TCFC film. The slope of the M-H loop decreases and leads to a larger anisotropy field, as it can be observed in Fig. 7.15 (b). Thus, under a specific magnetic field that is lower than the anisotropy field, the magnetization rises with tensile strain and reduces with compressive strain. The ideal M-E curves of TCFC/PMN-PZT/Au heterostructure should be in the inverted-butterfly shape if TCFC film is perfectly uniform and without defects.

It can be noticed that the M-E curves are in the inverted-butterfly shape when the magnetic field is higher than 800 Oe in Fig. 7.13, which is in accordance with the ideal deduction. However, here are some issues that the ideal model cannot explain: a). The double magnetization peaks around EC; b). The appearance of direct butterfly-like M-E curves when the magnetic field is lower than 800 Oe; c). The inverse process of the butterfly M-E curves.

As it is mentioned in Chapter 4, the surface of ceramic substrates is uneven, and many holes present. The diameter of the holes can be tens of micrometers, while the thickness of TCFC film is only 200 nm. Some parts of the TCFC film can be deposited into the holes of ceramic substrates. Thus the film is not exactly on the same plane. The magnetic anisotropy of different parts of TCFC film on different planes can lead to a complicated and mixed effect. As a result, the TCFC film is not perfect; it exhibits large remnant magnetization and non-zero HC along the hard axis. The hysteresis part, which appears when the magnetic field is lower than 800 Oe in Fig. 7.15, is the reason for the issues listed above. Since ceramic is polycrystalline, it consists of many different oriented grains. It is challenging to bring out a model to simulate the ME coupling effect in this complex system. However, a hypothesis is given above to explain the wired phenomena in TCFC/PMN-PZT/Au heterostructure.

The ∆M and CME coefficient is calculated and given in Fig. The maximum of the CME coefficient under the specific magnetic field can be reached around the EC of PMN-PZT. Figure 7.17 shows the value of the αmax versus bias magnetic field. With the bias magnetic field increasing, the αmax firstly decreases and then increases. The maximum of αmax is reached under a bias magnetic field of 300 Oe with a significant value of 136.6 ×10 -8 s/m. At 300 Oe, the unfinished inverse of the M-E curve at EC of PMN-PZT accounts for the significant CME coefficient. As the bias magnetic field further rises, the αmax reduces quickly and is lower than 50 ×10 -8 s/m until H = 2800 Oe. It is necessary to point out that, the αmax can reach 56.2 ×10 -8 s/m with an absence of the magnetic field. The slope of M-H loops changes implies the transformation of magnetic anisotropy. In Fig. 20 (b), the slope vs. E curve is in the similar shape of S-E curves of PMN-PZT, once again revealing the strain-mediated magnetic anisotropy transformation along the OOP direction. When there is a tensile OOP strain generated in PMN-PZT, the slope of M-H loops increases, indicating that the anisotropy axis tilts to OOP direction. While a compressive OOP strain is generated, the slope reduces, and the anisotropy axis deviates from the OOP direction.

CME Effect along Easy Axis

CME Effect along OOP Direction

As a result of E-induced anisotropy transformation, the E-field controlled magnetization variations can be observed along the OOP direction.

Summary

In this part of the work, the [(TbCo2)/(FeCo)]20/PMN-PZT/Au heterostructure has been prepared, and the E-field induced magnetism along different axes has been studied.

The M-H loops under different electric fields and E-field controlled magnetization variations along the hard axis of TCFC film have been investigated. The M-E curves are in accord with the butterfly S-E curve of the PMN-PZT substrate. An inverse process of M-E curves happens with the bias magnetic field increasing, and double magnetization peaks occur around the EC of PMN-PZT substrate. The appearances of the inverse transition of M-E curves and double peaks in this heterostructure may be due to the uneven distribution of TCFC film on a ceramic substrate with inevitable holes. The maximum of αCME is at a substantial value of 136.6 ×10 -8 s/m with a bias magnetic field of 300 Oe. The weird inverse process of M-E curves has provided a possibility to reach a substantial CME coupling effect. Along the easy axis of TCFC film, the E-induced magnetization variation is suppressed by the highly intrinsic crystallographic orientation with a magnetic field. Along OOP direction, clear evidence of E-field induced magnetic anisotropy transformation has been observed, i.e., the slope shifts of M-H loops with the electric field. The curve of slope versus electric field agrees with the strain evolution of PMN-PZT ceramic under an electric field. This is the direct evidence of the occurrence of strain-mediated magnetic anisotropy transformation in TCFC/PMN-PZT/Au heterostructure. The E-field controlled magnetization variations and CME coefficient along the OOP direction have also been demonstrated.

Conclusions

The rare-earth Tb modified 3d-transition alloys (TbCo2)/(FeCo) is used to couple with PLZST and PMN-PZT ceramic substrates in this part. The 1. Along the easy axis of TCFC film, the E-induced magnetization variation is suppressed by the highly intrinsic crystallographic orientation with the magnetic field.

2. Along the hard axis, a reverse process of M-E curves appears with the bias magnetic field increasing. The appearances of the inverse transition of M-E curves may be due to the uneven distribution of TCFC film on a ceramic substrate with inevitable holes. In TCFC/PLZST/Au, the maximum of ∆M/M0 can reach -9.8% with an absence of a magnetic field, and the αCME is at the negative maximum of -10.3×10 -8 s/m at the same time. The positive maximum of αCME is at 500 Oe with a value of 12.7×10 -8 s/m.

In TCFC/PMN-PZT/Au, the maximum of αCME can reach a large value of 136.6 ×10 -8 s/m with a bias magnetic field of 300 Oe.

3. Along OOP direction, the E-induced magnetism is not visible in TCFC/PLZST/Au heterostructure, which can be attributed to that the energy of Einduced strain in PLZST is not enough to deflect the magnetic moment from in-plane to OOP direction. However, in TCFC/PMN-PZT/Au heterostructure, clear evidence of E-field induced magnetic anisotropy transformation has been observed, i.e., the slope shifts of M-H loops with the electric field. The curve of slope versus electric field agrees with the strain evolution of PMN-PZT ceramic under an electric field.

YIG/AFE and YIG/FE Heterostructures in Full Thin-Film Form

Extensive investigations have been devoted to the research of multiferroic composites in full thin-film in past years with the development of integration and miniaturization. [START_REF] More-Chevalier | Magnetoelectric coupling in Pb(Zr,Ti)O3-Galfenol thin film heterostructures[END_REF][START_REF] Ma | Recent progress in multiferroic magnetoelectric composites: from bulk to thin films[END_REF][START_REF] Cibert | Magnetoelectric coupling in Tb0.3Dy0.7Fe2/Pt/PbZr0.56Ti0.44O3 thin films deposited on Pt/TiO2/SiO2/Si substrate[END_REF][START_REF] Hu | Multiferroic magnetoelectric nanostructures for novel device applications[END_REF][START_REF] Martin | Multiferroics and magnetoelectrics: thin films and nanostructures[END_REF][START_REF] Ortega | Multifunctional magnetoelectric materials for device applications[END_REF] Compared with bulk ME composites, multiferroic thin films have unique superiorities. The multiferroic thin films show great potential in integrated magnetic/electric devices, such as sensors, microelectromechanical system, high-density memories, and spintronics. [START_REF] Hu | Multiferroic magnetoelectric nanostructures for novel device applications[END_REF][START_REF] Ortega | Multifunctional magnetoelectric materials for device applications[END_REF][START_REF] Deng | Magnetoelectric effect in multiferroic heteroepitaxial BaTiO3-NiFe2O4 composite thin films[END_REF][START_REF] Wan | Magnetoelectric CoFe2O4-Pb(Zr,Ti)O3 composite thin films derived by a sol-gel process[END_REF][START_REF] Cibert | Magnetoelectric coupling in Tb0.3Dy0.7Fe2/Pt/PbZr0.56Ti0.44O3 thin films deposited on Pt/TiO2/SiO2/Si substrate[END_REF] In the previous parts, the CME coupling effect has been demonstrated in AFE-and FE-based bulk heterostructures. The large thickness of AFE ceramic substrates requires an application of high voltage to reach the switching field, which is a big challenge for future applications. Thus, it is necessary to develop AFE-and FE-based multiferroic films and compare their performance in thin-film form. The ferrite YIG films are used to deposit on AFE and FE films for their good coupling effect.

Preparation of YIG/AFE and YIG/FE Heterostructures in Full-

Thin-Film Form

In these heterostructures, the PbZrO3 (PZ) AFE film Pb(Zr, Ti)O3 (PZT) FE film were grown on SiO2/Si substrate where the LaNiO3 (LNO) film was introduced as the bottom electrodes. Then Pt top electrodes and the following YIG films were deposited.

The YIG/Pt/PZ/LNO/SiO2/Si and YIG/Pt/PZT/LNO/SiO2/Si heterostructure in full thin-film form can be prepared.

LNO film Growth

The LNO electrode layer was prepared by using an "Addax" RF sputtering system with an LNO ceramic target. The sputtering condition and annealing process of LNO films have been explored by several researchers and well known in the group. [START_REF] Li | Influence of LNO top electrodes on electrical properties of KNN/LNO thin films prepared by RF magnetron sputtering[END_REF][START_REF] Sama | On the influence of the top and bottom electrodes-A comparative study between Pt and LNO electrodes for PZT thin films[END_REF][START_REF] Ge | Effect of electrode materials on the scaling behavior of energy density in Pb(Zr0.96Ti0.03)Nb0.01O3 antiferroelectric films[END_REF][START_REF] Li | Characteristics of highly (001) oriented (K,Na)NbO3 films grown on LaNiO3 bottom electrodes by RF magnetron sputtering[END_REF][START_REF] Yang | Orientation control of LaNiO3 thin films by RF magnetron sputtering with different oxygen partial pressure[END_REF] The details of the sputtering condition are given in Table 8.1. After the deposition, the samples were annealed in a furnace to be crystallized.

The annealing process is given in Table 8.2. Cooling Rate (℃/min) 1

Atmosphere Air Figure 8.1 shows the XRD patterns of LNO films on SiO2/Si substrates, and it can be observed that the LNO films can be well crystallized on SiO2/Si substrates at 750 ℃.

The LNO film is highly (100)-oriented, which will favor the (100)-oriented films on the top. hour. The PZ layer was spin-coated and then pyrolyzed at 350 °C and at 600 °C respectively. These steps were repeated for six times to achieve the thickness of 500 nm. Finally, the films were annealed at 750 °C in Rapid Temperature Annealing.

Usually the annealing temperature of PZ films is 700 °C, here a high annealing temperature of 750 °C is to adapt to the following annealing process of YIG films.

The PZ film with a thickness of 500 nm shows good crystallization on LNO/ SiO2/Si substrates, as given in Fig. 8.2. The highly (100)-oriented PZ films are formed with no impurity phases detected.

Preparation of Pt/PZT/LNO/SiO2/Si ME Heterostructures

The PZT film with a thickness of 500 nm was prepared by using an "Addax" RF sputtering system with a powder target in a composition of PbZr0.53Ti0.47O3. The target was fabricated with an excess of 20% Pb to compensate for the deficient of Pb during the deposition and annealing process. The sputtering condition and the annealing process of PZT films have been explored in the group. [START_REF] Sama | On the influence of the top and bottom electrodes-A comparative study between Pt and LNO electrodes for PZT thin films[END_REF][START_REF] Herdier | A silicon cantilever beam structure for the evaluation of d31, d33 and e31 piezoelectric coefficients of PZT thin films[END_REF][START_REF] Sama | Bottom and top electrodes nature and PZT film thickness influence on electrical properties[END_REF] The details of the sputtering condition are given in Table 8.3. After the deposition, the samples were annealed in a furnace to be crystallized.

The annealing process is given in Table 8.4. Usually, the annealing temperature of PZT films is 650°C. In this work, we have used 750°C for the simple reason that the YIG films must be annealing at 750°C. We have systematically verified that the structure of PZT is not affected by this high annealing temperature, and in particular no second phase (lead deficient) appears, as given in Figure 8.5. The AFM surface morphology of YIG films on SiO2/Si substrate in Figure 8.12 (a). A dense granular-like surface was observed. The magnetic domains of YIG films on SiO2/Si substrate were also locally probed by MFM technique. As shown in Figure 8.12 (b), where a maze shape is obtained for domains, with yellow/clear and blue/dark contrasts corresponding to upward (repulsive interaction) and downward (attractive interaction) OOP magnetization, respectively. Such complex irregular domain pattern was already explained by the similar magnetic anisotropy and magnetostatic energies in the film. [START_REF] More-Chevalier | Reversible control of magnetic domains in a Tb0.3Dy0.7Fe2/Pt/PbZr0.56Ti0.44O3 thin film heterostructure deposited on Pt/TiO2/SiO2/Si substrate[END_REF][START_REF] Mudivarthi | Magnetic domain observations in Fe-Ga alloys[END_REF][START_REF] Wang | The correlation between the first magnetocrystalline anisotropy constant K1 and the composition of giant magnetostrictive alloy Tb-Dy-Fe[END_REF] Besides, no-correlation is observed between AFM and MFM images confirming the contrasts seen on MFM image are exclusively due to the magnetic interaction. 1). The OOP magnetic response of YIG films are too weak to be detected by MFM system. As we have demonstrated in Chapter 6, the polycrystalline YIG films present an "easy plane" in the plane, and the OOP magnetization can reach the saturation under the application of a very high magnetic field. In Chapter 6.2, the in-plane saturation magnetic field of YIG films is 200 Oe, while the OOP saturation magnetic field is 1500

Oe. Thus, the magnetic response of YIG film is too weak along OOP direction.

2). During the MFM measurements, a proper distance between the MFM tip and the surface of YIG films is necessary to reach a good resolution of AFM image and MFM image. Since the OOP magnetic response of YIG films is weak, a very low distance is needed to obtain the magnetic response. However, if the MFM tip is too close to the surface of YIG films, the morphology signal would be strong enough to cover up the magnetic signal. A low distance of 15 nm and 10 nm has been tried, but we couldn't obtain distinct magnetic responses.

3). The quality of PZ films has an influence on the quality of YIG films. The multistep preparation process may lead to some side effect, especially the multiple annealing process have an effect on PZ film. The Pb is a volatile constituent, and the volatilization happens at high temperature. The volatilization process not only lead to Pb loss in PZ film, but also cause some Pb diffusion into other layers, even though it is not detected by XRD system.

Even with the application of electric field, the variations of magnetic response is still a challenge. As it has been presented in Chapter 6, the E-field induced magnetization is very weak along OOP direction.

YIG/Pt/PZT/LNO/SiO2/Si Heterostructure

The characterization of MFM image in YIG/Pt/PZT/LNO/SiO2/Si heterostructure has also met the same challenge as YIG/Pt/PZ/LNO/SiO2/Si.

Conclusions

The ME composites in thin-film form were fabricated in this part -YIG/Pt/PZ/LNO/SiO2/Si and YIG/Pt/PZT/LNO/SiO2/Si heterostructures. The PZ and PZT films have shown good electrical properties, but the MFM image cannot be successfully characterized in this case. The in-plane anisotropy field of YIG films has constrained the detectable magnetic signal through MFM system along OOP direction.

The multiple annealing process of PZ and PZT films may also has a side effect because of Pb volatilization. In addition, the CME coupling effect of polycrystalline YIG is weak along OOP direction, which would also be a challenge for future measurements.

Therefore, this part of work has pointed out the way to fabricate CME composites in thin-film form. A good choice of ferromagnetic phases should be the films with a strong magnetic response along OOP direction and low annealing temperature or no annealing process needed, such as 3d-transition metals (Fe, Co, and Ni) and alloys (Terfenol-D, Galfenol).

and the detailed characterizations have been performed. The well-polished substrates can reach a low average roughness (< 20 nm) despite the inevitable defects in ceramics.

The EC of PMN-PZT ceramics is as low as 4.2 kV/cm, while the E-induced in-plane and out-of-plane strain can reach -0.27% and 0.08% at EC, respectively. The AFE-FE switching field of PLZST is 5 kV/mm with a maximum strain of 0.2 %.

Second, the formation of NMG/PLZST heterostructure has been performed, and the E-field controlled magnetization switching in an AFE-based multiferroic composite has been firstly reported. The evolution of (ΔM/M0) as a function of E is in good agreement with the ceramic strain evolution and secondly that the maximum ΔM/M0 variation, which attains a maximum value of 15%, is reached at 0 Oe. Even the CME coefficient is small due to the bad quality of NMG film on PLZST ceramic, and these results reveal the great potential of AFE materials in the CME coupling effect for the first time.

Third, we investigated the E-field control of magnetism in YIG/Pt/PLZST/Pt, and YIG/Pt/PMN-PZT/Pt heterostructures, wherein the in-plane and out-of-plane CME coupling effects were explored. In these two heterostructures, the E-field induced strain is transferred to YIG film and leads to a magnetic anisotropy transformation, which results in the manipulation of magnetism. Because of the different strain evolution of PLZST and PMN-PZT substrates under the electric field, the CME coupling effect in YIG/Pt/PLZST/Pt and YIG/Pt/PMN-PZT/Pt heterostructures shows some differences.

The considerable E-induced phase transition strain in PLZST substrate leads to maximum in-plane αCME (11.6 × 10 -8 s/m) at 0 Oe and a maximum of in-plane ∆χ/χ0 with an absence of the magnetic field, and the αCME is at the negative maximum of -10.3×10 -8 s/m at the same time. The positive maximum of αCME is at 500 Oe with a value of 12.7×10 -8 s/m. In TCFC/PMN-PZT/Au, the maximum of αCME can reach a significant value of 136.6 ×10 -8 s/m with a bias magnetic field of 300 Oe. Along OOP direction, the E-field induced magnetism is not visible in TCFC/PLZST/Au heterostructure, which can be attributed to that the energy of E-field induced strain in PLZST is not enough to deflect the magnetic moment from in-plane to OOP direction.

However, in TCFC/PMN-PZT/Au heterostructure, clear evidence of E-field induced magnetic anisotropy transformation has been observed, i.e., the slope shifts of M-H loops with the electric field. The curve of the slope versus electric field agrees with the strain evolution of PMN-PZT ceramic under the electric field.

Finally, we have explored the YIG/AFE and YIG/FE multiferroic heterostructure in thin-film form. The results indicate that a ferromagnetic phase with a strong magnetic response along OOP direction is needed for the MFM characterization to evaluate the ME performance of these structures.

Future Work

Despite the progress made in this dissertation, further work is needed to understand the AFE-and FE-based multiferroic heterostructures fully.

The following future work has been proposed:

1. A deeper study of TCFC/PLZST and TCFC/PMN-PZT heterostructures is needed. The theoretical simulation and local characterizations of magnetic response under certain magnetic and electric fields are necessary for us to understand the inverse of M-E curves.

2. In this thesis, the E-field induced magnetization switching and E-field tunable magnetic susceptibility have been demonstrated in different heterostructures. However, the E-field induced FMR shifts have not been done to compare the differences of CME coupling effect in AFE-and FE-based multiferroic heterostructures at high frequency.

1 27 Table 5 . 1 46 Table 5 . 2 46 Table 6 . 1 55 Table 6 . 2 56 Table 6 . 3 60 Table 6 . 4 66 Table 6 . 5 70 Table 7 . 1 95 Table 8 . 1 111 Table 8 . 2 112 Table 8 . 3 115 Table 8 . 4 115 Table 8 . 5 Figure 2 . 1 5 Figure 2 . 2 6 Figure 2 . 3 7 Figure 2 . 4 8 Figure 2 . 5 10 Figure 2 . 6 11 Figure 2 . 7 11 Figure 2 . 8 12 Figure 2 . 9 15 Figure 2 . 10 16 Figure 2 . 11 19 Figure 2 . 21 Figure 2 . 13 22 Figure 2 . 14

 12751465246615562566360646665707195811118211283115841158521522623724825102611271128122915210162111922121322214 Figure 2.1 Schematic of the four primary ferroics.[42] ......................................... 5 Figure 2.2 (a) Hysteresis loop, and (b) longitudinal S-E loop of FEs. [45] ........... 6 Figure 2.3 (a) A typical ABO3 united-cell structure, and (b) Primitive cell schematic diagram of antiferroelectric (tetragonal structure) and ferroelectric (hexagonal structure).[46] .............................................................................. 7 Figure 2.4 (a) Double hysteresis loop, and (b) longitudinal S-E loop of AFEs.[47] ........................................................................................................................ 8 Figure 2.5 The classification of magnetic materials. ........................................... 10 Figure 2.6 The magnetic hysteresis loop of FM materials. ................................. 11 Figure 2.7 The magnetic domain of ferromagnetic materials. ............................. 11 Figure 2.8 Schematic illustrating magnetic-elastic-electric couplings in multiferroic materials.[56] ............................................................................ 12 Figure 2.9 Schematic illustration of three bulk composites with the three common connectivity schemes: (a) 0-3 particulate composite, (b) 2-2 laminate composite, and (c) 1-3 fiber/rod composite.[61] ......................................... 15 Figure 2.10 Schematic illustration of strain-mediated ME effect in a composite system consisting of a magnetic layer (purple) and ferroelectric layer (pink): (a) DME effect (b) CME effect.[78] ............................................................. 16 Figure 2.11 Magnetization as functions of temperature for different values of the electric (E) fields.[98] ................................................................................... 19 Figure 2.12 Heat-assisted strain-induced magnetization switching. (a) Heating and cooling of multiferroic layers during and after a laser pulse. The inset shows the temperature profile in the free layer; (b,c) heat-assisted magnetization switching starting from states "1" and "0" with two different stress directions.[23] ............................................................................................... 21 Figure 2.13 Induced magnetic moment variance with (a) the bias magnetic field, (b) the bias electric field, (c) the angle θ between the sample plane and magnetic field, and (d) the thickness ratio between Terfenol-D and PLZST at a resonance frequency of 365 kHz.[27] ........................................................ 22 Figure 2.14 The magneto-optical Kerr effect signal converted to the oscillation amplitude of the magnetization vector (∆γ) measured at different AC voltages.[99] ................................................................................................. 22

Figure 2 . 23 Figure 2 . 16

 223216 Figure 2.15 (a) M-H hysteresis loops of the austenite and martensite NMG thin films, (b) Out-of-plane P-E hysteresis loops of the PLZST films with austenite and martensite NMG film on the top, before and after annealing, respectively.[22] ........................................................................................... 23 Figure 2.16 Enlarged view of PLZT (001) peaks measured under various tilt angles for (a) the as-deposited, and (b) the annealed samples, respectively. (c) dspacing of PLZT (001) plane as a function of sin2ψ for the two samples.[37] ...................................................................................................................... 23

Figure 2 . 24 Figure 3 . 1 A 29 Figure 3 . 2

 224312932 Figure 2.17 (a) The schematic of FeGaB film deposited on the top or on the side of PSZT ceramics, (b) P-E loop of PZST ceramic, (c) FMR spectra under varying E-field of FeGaB(top)/PSZT composites and inset the FMR field dependence of E-field, and (d) magnetization switch of FeGaB(top)/PSZT at 40.5 Oe induced by voltage impulse.[26] ..................................................... 24 Figure 3.1 A principle schematic representation of RF magnetron sputtering[103] ...................................................................................................................... 29 Figure 3.2 RF magnetron sputtering system of Plassys. ...................................... 30

Figure 3 . 3

 33 Figure 3.3 (a) Home-made VSM system, (b) Connected sample. ....................... 33

Figure 4 . 1

 41 Figure 4.1 Piezoelectric coefficient d33, electromechanical coupling factor kp, and quality factor Qm of the x PMN-(1-x) PZT ceramics.[115] ........................ 35

Figure 4 . 2

 42 Figure 4.2 (a) AFM picture of the polished PMN-PZT substrate surface, (b) XRD patterns of PMN-PZT ceramic powders. ...................................................... 36

Figure 4 . 3 37 Figure 4 . 4

 433744 Figure 4.3 SEM picture of the polished PMN-PZT substrate surface. ................ 37 Figure 4.4 Schematic of PMN-PZT ceramic substrates for electrical measurements. ...................................................................................................................... 37

Figure 4 . 5

 45 Figure 4.5 The P-E loop of PMN-PZT ceramic substrate. .................................. 38

Figure 4 . 6

 46 Figure 4.6 (a) The in-plane strain and (b) out-of-plane strain of PMN-PZT ceramic substrate. ....................................................................................................... 38

Figure 4 . 7

 47 Figure 4.7 Ternary phase diagram of the Pb0.97La0.02(Zr, Sn, Ti)O3 system.[121] ...................................................................................................................... 40

Figure 4 . 8 41 Figure 4 . 9

 484149 Figure 4.8 (a) The AFM picture and (b) XRD pattern of PLZST ceramic substrate. ...................................................................................................................... 41 Figure 4.9 The surface SEM picture of the PLZST ceramic substrate. ............... 41

Figure 4 . 42 Figure 5 . 1

 44251 Figure 4.10 The P-E loop and S-E curve of the PLZST ceramic substrate. ........ 42 Figure 5.1 The XRD patterns of NMG films on Si substrates. ............................ 47

Figure 5 . 2 48 Figure 5 . 3 48 Figure 5 . 4 49 Figure 5 . 5 50 Figure 5 . 6

 524853485449555056 Figure 5.2 The magnetic hysteresis loops of NMG films on Si substrates annealed at different temperatures. .............................................................................. 48 Figure 5.3 Schematic of NMG/PLZST/NMG multiferroic heterostructure. ....... 48 Figure 5.4 (a) The cross-sectional SEM picture of NMG/PLZST and (b)XRD patterns of NMG/Si, PLZST, and NMG/PLZST samples. ........................... 49 Figure 5.5 The M-H loop of NMG/PLZST/NMG heterostructure. ..................... 50 Figure 5.6 The M-H loop of NMG/PLZST/NMG heterostructure with and without an electric field. ............................................................................................ 50

Figure 5 . 7 51 Figure 5 . 8

 575158 Figure 5.7 (a) (ΔM/M0)-E curves under bias magnetic field of 0, 100 and 500 Oe, and (b) S-E curve and |ΔM/M0|-E loop measured at 500 Oe of NMG/PLZST/NMG heterostructure............................................................. 51 Figure 5.8 The (ΔMmax/M0)-H curve of NMG/PLZST/NMG heterostructure. .... 52

Figure 5 . 9

 59 Figure 5.9 (a) ΔM/M0 and CME coefficient versus E under a bias magnetic field of 500 Oe, (b) the maximum of CME coefficient versus bias H under positive and negative voltages of NMG/PLZST/NMG sample. ................................ 53

Figure 6 . 1 57 Figure 6 . 2 58 Figure 6 . 3 58 Figure 6 . 4 59 Figure 6 . 5

 615762586358645965 Figure 6.1 (a) The XRD patterns and (b) M-H loops of YIG films on SiO2/Si substrates with different annealing temperatures. ........................................ 57 Figure 6.2 Schematic of YIG/Pt/PLZST (PMN-PZT)/Pt heterostructure (a) front side and (b) backside. ................................................................................... 58 Figure 6.3 Schematic of (a) YIG/Pt/PLZST/Pt and (b) YIG/Pt/PMN-PZT/Pt ME heterostructures. ............................................................................................ 58 Figure 6.4 The XRD pattern and inset cross-sectional SEM image of (a) YIG/Pt/PLZST/Pt and (b) YIG/Pt/PMN-PZT/Pt heterostructures. .............. 59 Figure 6.5 The in-plane M-H loops of YIG/Pt/PLZST/Pt heterostructure. ......... 60

Figure 6 . 6

 66 Figure 6.6 (a) M-H loops, inset the enlarged M-H loops, (b) HC-E curve of YIG/Pt/PLZST/Pt heterostructure. ............................................................... 61

Figure 6 . 7 61 Figure 6 . 8 63 Figure 6 . 9

 6761686369 Figure 6.7 (a) χ-H curves with different electric fields and (b) χ-E curve at 10 Oe of YIG/Pt/PLZST/Pt heterostructure. ........................................................... 61 Figure 6.8 The M-E curves of YIG/Pt/PLZST/Pt at 0 Oe, 15 Oe, and 30 Oe. The labeled M0i and M0f are the initial magnetization and final magnetization at 0 kV/cm. .......................................................................................................... 63 Figure 6.9 (a) The relative magnetization change versus magnetic field (∆M/M0f -H) curve, (b) (M0i-M0f)/M0f -H curve of YIG/Pt/PLZST/Pt heterostructure. ...................................................................................................................... 64

Figure 6 . 10 64 Figure 6 . 65 Figure 6 . 12 66 Figure 6 . 13 67 Figure 6 . 14

 610646656126661367614 Figure 6.10 The magnetization versus electric field curves of YIG/Pt/PLZST/Pt (a) 0 Oe, (b) 30 Oe. ....................................................................................... 64 Figure 6.11 ∆M/M0-E and αCME-E curves at 0 Oe of YIG/Pt/PLZST/Pt heterostructure. ............................................................................................. 65 Figure 6.12 The CME coefficient versus bias magnetic field curves of YIG/Pt/PLZST/Pt heterostructure. ............................................................... 66 Figure 6.13 The out-of-plane M-H loops of YIG/Pt/PLZST/Pt heterostructure with different electric fields. ................................................................................. 67 Figure 6.14 The out-of-plane M-E curves of YIG/Pt/PLZST/Pt heterostructure (a) 500 Oe, (b) 1000 Oe. .................................................................................... 68

Figure 6 . 68 Figure 6 . 16 69 Figure 6 . 17

 66861669617 Figure 6.15 (a) The in-plane M-E curve and (b) out-of-plane M-E curve of YIG/Pt/PLZST/Pt heterostructure. ............................................................... 68 Figure 6.16 Schematic of the veer of the anisotropy axis in YIG film. ............... 69 Figure 6.17 The normal and inset the enlarged M-H loop of YIG/Pt/PMN-PZT/Pt heterostructure. ............................................................................................. 70

Figure 6 .

 6 Figure 6.18 (a) M-H loops with different electric fields, inset the enlarged M-H loops, (b) Mr-E and HC-E curves of YIG/Pt/PMN-PZT/Pt heterostructure. 71

Figure 6 .Figure 6 . 20 Figure 6 . 21

 6620621 Figure 6.19 (a) χ-H curves with different electric fields and (b) χ-E curve at 15, 20, and 25 Oe of YIG/Pt/PMN-PZT/Pt heterostructure. .................................... 72 Figure 6.20 The first and second cycle of M-E curves at 0 Oe in YIG/Pt/PMN-PZT/Pt heterostructure. ................................................................................. 73 Figure 6.21 The electric-field-dependent magnetization (M-E) in YIG/Pt/PMN-PZT/Pt under the bias magnetic field of 0, 10, 25, 50, 75, 100, 200 and 500 Oe ................................................................................................................. 74

Figure 6 . 75 Figure 6 . 23

 675623 Figure 6.22 (a) ∆M/M0-E and αCME-E curves at 0 Oe, and (b) |∆M/M0| -H and (αCME) -H curves of YIG/Pt/PMN-PZT/Pt heterostructure. .......................... 75 Figure 6.23 The out-of-plane strain evolution of PMN-PZT ceramic substrate. . 76

Figure 6 . 77 Figure 6 .

 6776 Figure 6.24 M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure under cycling electric field with different amplitudes at 0 Oe. (a) M-E curves at 1.6, 3.0, 4.0, 4.6 kV/cm, (b) 4.6, 5.2 6.0, 7.0 kV/cm, (c) 7.0, 9.0, 12 kV/cm. .................. 77 Figure 6.25 M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure under cycling electric field with different amplitudes at 0 Oe. (a) Poled at +12 kV/cm and (b) poled at -12 kV/cm.................................................................................. 78

Figure 6 . 78 Figure 6 . 79 Figure 6 . 28 80 Figure 6 . 29 81 Figure 6 . 30 82 Figure 6 . 31 83 Figure 6 . 32

 67867962880629816308263183632 Figure 6.26 χ-E curves of YIG/Pt/PMN-PZT/Pt heterostructure under cycling electric field with different amplitudes at 20 Oe. ......................................... 78 Figure 6.27 (a) The M-E curve with an electric field cycle from -4.6 kV/cm to 4.6 kV/cm, (b) the time-dependent magnetization switched by the electric field at 0 Oe of YIG/Pt/PMN-PZT/Pt heterostructure. ............................................. 79 Figure 6.28 The χ-E curve with an electric field cycle from -4 kV/cm to 4 kV/cm at 20 Oe of YIG/Pt/PMN-PZT/Pt heterostructure. ....................................... 80 Figure 6.29 The electric-field dependent M-H loops of YIG/Pt/PMN-PZT/Pt heterostructure. ............................................................................................. 81 Figure 6.30 The out-of-plane M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure. ...................................................................................................................... 82 Figure 6.31 The in-plane M-E curve at 0 Oe and out-of-plane M-E curve at 500 Oe of YIG/Pt/PMN-PZT/Pt heterostructure. ................................................ 83 Figure 6.32 Schematic of electric field induced anisotropy change in YIG/Pt/PMN-PZT/Pt heterostructure. .......................................................... 83

Figure 7 . 1

 71 Figure 7.1 (a) Schematic of TCFC multilayers and (b) TCFC/PLZST/Au heterostructure. ............................................................................................. 88

Figure 7 . 2 89 Figure 7 . 3 90 Figure 7 . 4 91 Figure 7 . 5 93 Figure 7 . 6

 728973907491759376 Figure 7.2 (a) Front side of TCFC/PLZST (PMN-PZT)/Au heterostructure, (b) picture of TCFC/PLZST (PMN-PZT)/Au connected to sample holder, (c) front side and (d) backside of TCFC/PLZST (PMN-PZT)/Au connected with wires. ............................................................................................................ 89 Figure 7.3 The M-H loops of TCFC/PLZST/Au heterostructure along with different directions. ....................................................................................... 90 Figure 7.4 The M-H loop of TCFC/PLZST/Au heterostructure along the hard axis. ...................................................................................................................... 91 Figure 7.5 The M-E curves of TCFC/PLZST/Au heterostructure along the hard axis at 0, 50, 100, 190, 500, 1000, and 1500 Oe. ......................................... 93 Figure 7.6 The ∆M-E, and αCME-E curves of TCFC/PLZST/Au heterostructure along the hard axis at 0, 50, 100, 190, 500, 1000, and 1500 Oe................... 94

Figure 7 . 7 95 Figure 7 . 8

 779578 Figure 7.7 (a) The ∆M/M0-H curve and (b) αCME-H curve of TCFC/PLZST/Au heterostructure along hard axis. .................................................................... 95 Figure 7.8 The M -H loops along the easy axis of TCFC/PLZST/Au heterostructure under different electric fields. .............................................. 96

Figure 7 . 9 97 Figure 7 . 10 98 Figure 7 . 11 99 Figure 7 . 12 100 Figure 7 . 13 101 Figure 7 . 14 102 Figure 7 . 103 Figure 7 . 16 104 Figure 7 . 17 105 Figure 7 . 18 106 Figure 7 . 106 Figure 7 . 107 Figure 7 . 21 108 Figure 7 . 22 108 Figure 8 . 1 113 Figure 8 . 2 113 Figure 8 . 3

 79977109871199712100713101714102710371610471710571810671067107721108722108811138211383 Figure 7.9 The M -H loops along the OOP direction of TCFC/PLZST/Au heterostructure under different electric fields. .............................................. 97 Figure 7.10 The M-H loops of TCFC/PMN-PZT/Au heterostructure along with different directions. ....................................................................................... 98 Figure 7.11 The M-H loops along the hard axis of TCFC/PMN-PZT/Au heterostructure under different electric fields. .............................................. 99 Figure 7.12 The M-E curves along the hard axis of TCFC/PMN-PZT/Au heterostructure at 0 and 40 Oe. ................................................................... 100 Figure 7.13 The M-E curves along the hard axis of TCFC/PMN-PZT/Au heterostructure under different magnetic fields. ......................................... 101 Figure 7.14 The ideal M-H loops along the hard axis of a single crystal with different strains. .......................................................................................... 102 Figure 7.15 The M-H loops along the hard axis of TCFC/PMN-PZT/Au heterostructure under different magnetic fields. ......................................... 103 Figure 7.16 The ∆M-E and αCME-E curves of TCFC/PMN-PZT/Au heterostructure along the hard axis at 0, 40, 80, 160, 300, 500, 800, 1000, 1300, 1500 and 2800 Oe. ............................................................................ 104 Figure 7.17 The 𝛼𝑚𝑎𝑥 -H curve along the hard axis of TCFC/PMN-PZT/Au heterostructure. ........................................................................................... 105 Figure 7.18 The M-H loops along the easy axis of TCFC/PMN-PZT/Au heterostructure under different electric fields. ............................................ 106 Figure 7.19 The M-H loops along the OOP direction of TCFC/PMN-PZT/Au heterostructure under different electric fields. ............................................ 106 Figure 7.20 The M-H loops along the easy axis of TCFC/PMN-PZT/Au heterostructure under different electric fields. ............................................ 107 Figure 7.21 The M-E curves along the OOP direction of TCFC/PMN-PZT/Au heterostructure (a) at 5000 Oe and (b) at 15000 Oe. .................................. 108 Figure 7.22 The ∆M/M0-E and αCME-E curves along the OOP direction of TCFC/PMN-PZT/Au heterostructure (a) at 5000 Oe and (b) at 15000 Oe. .................................................................................................................... 108 Figure 8.1 The XRD patterns of LNO films on SiO2/Si substrates. .................. 113 Figure 8.2 The XRD patterns of PZ films on LNO/SiO2/Si substrates. ............ 113 Figure 8.3 Schematic of Pt/PZ/LNO/SiO2/Si heterostructure. .......................... 114

Figure 8 . 4 114 Figure 8 . 5 116 Figure 8 . 6 117 Figure 8 . 7 117 Figure 8 . 8 118 Figure 8 . 9 119 Figure 8 . 120 Figure 8 .

 84114851168611787117881188911981208 Figure 8.4 (a) The P-E loops and (b) εr-E and tanδ-E curves of Pt/PZ/LNO/SiO2/Si heterostructure. ........................................................................................... 114 Figure 8.5 Schematic of Pt/PZT/LNO/SiO2/Si heterostructure. ........................ 116 Figure 8.6 (a) The P-E loops and (b) εr-E and tanδ-E curves of Pt/PZT/LNO/SiO2/Si heterostructure. ........................................................ 117 Figure 8.7 Schematic of Pt/PZ/LNO/SiO2/Si heterostructure. .......................... 117 Figure 8.8 XRD patterns of (a) YIG/Pt/PZ/LNO/SiO2/Si and (b) YIG/Pt/PZT/LNO/SiO2/Si heterostructures. .............................................. 118 Figure 8.9 In-plane M-H loops of YIG/SiO2/Si and YIG/Pt/PZ/LNO/SiO2/Si heterostructures. .......................................................................................... 119Figure 8.10 (a) The MFM system, and (b) schematic of the YIG/Pt/PZ/LNO/SiO2/Si heterostructure with probes. ............................... 120Figure 8.11 (a) AFM surface morphology and (b) MFM image of YIG films on SiO2/Si substrate. ........................................................................................ 121

Chapter 4

 4 introduces the FE and AFE ceramic substrates prepared for CME coupling. The microstructure and electric properties of ceramic substrates are illustrated. Chapter 5 gives the first work of the E-field controlled magnetization variations in an AFE-based heterostructure. The NMG/AFE/NMG heterostructure is fabricated, and the CME coupling effect has been demonstrated. Chapter 6 and 7 present the CME coupling effect of FM films on FE ceramic substrates and AFE ceramic substrates, respectively. The multiferroic heterostructures of YIG/PLZST, YIG/PMN-PZT, [TbCo2/FeCo]20/PLZST and [TbCo2/FeCo]20/PMN-PZT are demonstrated in these two chapters. The E-field controlled magnetization switching and E-field tuning of the susceptibility of YIG film on PLZST and PMN-PZT ceramics are revealed, and the strain-induced anisotropy transformation of YIG film is also demonstrated. The CME coupling effect along different axes is presented in [TbCo2/FeCo]20 films on PLZST and PMN-PZT ceramics. In Chapter 8, the YIG/AFE and YIG/FE composites in thin-film form are fabricated, and the magnetic response of YIG films are tried to characterized. A summary of the thesis and several propositions of future work are provided in Chapter 9.

Figure 2 . 1

 21 Figure 2.1 Schematic of the four primary ferroics.[42]

Figure 2 . 2

 22 Figure 2.2 (a) Hysteresis loop, and (b) longitudinal S-E loop of FEs.[START_REF] Haertling | Ferroelectric ceramics: history and technology[END_REF]

Figure 2 . 3

 23 Figure 2.3 (a) A typical ABO3 united-cell structure, and (b) Primitive cell schematic diagram of antiferroelectric (tetragonal structure) and ferroelectric (hexagonal structure).[START_REF] Xu | Dielectric properties and field-induced phase switching of lead zirconate titanate stannate antiferroelectric thick films on silicon substrates[END_REF]

Figure 2 . 4

 24 Figure 2.4 (a) Double hysteresis loop, and (b) longitudinal S-E loop of AFEs.[START_REF] Tian | Progress of antiferroelectric pervoskite oxides[END_REF]

Figure 2 . 5

 25 Figure 2.5 The classification of magnetic materials.

Figure 2 . 6

 26 Figure 2.6 The magnetic hysteresis loop of FM materials.

Figure 2 . 8

 28 Figure 2.8 Schematic illustrating magnetic-elastic-electric couplings in multiferroic materials.[START_REF] Palneedi | Status and perspectives of multiferroic magnetoelectric composite materials and applications[END_REF]

Figure 2 . 9

 29 Figure 2.9 Schematic illustration of three bulk composites with the three common connectivity schemes: (a) 0-3 particulate composite, (b) 2-2 laminate composite, and(c) 1-3 fiber/rod composite.[START_REF] Nan | Multiferroic magnetoelectric composites: Historical perspective, status, and future directions[END_REF]

Figure 2 . 10

 210 Figure 2.10 Schematic illustration of strain-mediated ME effect in a composite system consisting of a magnetic layer (purple) and ferroelectric layer (pink): (a) DME effect (b) CME effect.[START_REF] Wang | Multiferroic magnetoelectric composite nanostructures[END_REF]

  Suchtelen first proposed the concept of the composite ME effect in 1972, the ME composites have gotten tremendous development. The appearance of the giant magnetostrictive rare-earth-iron alloy Terfenol-D has substantially simulated the upsurge of multiferroic ME composites research in the early 2000s. Since then, various multiferroic heterostructures have been intensively investigated to obtain the strong ME coupling effect.[86-92] Materials used for the ferromagnetic/ferrimagnetic phases include 3d-transition metals Fe, Co, and Ni and alloys such as ferrocobalt, Terfenol-D, Galfenol, spinel ferrites, hexagonal ferrites, and lanthanum manganites. The ferroelectric phases used in the previous studies were Pb(Zr, Ti)O3 (PZT), BaTiO3(BTO), lead magnesium niobate-lead titanate (PMN-PT), and poly (vinylidene fluoride) (PVDF).[START_REF] Lei | Straincontrolled magnetic domain wall propagation in hybrid piezoelectric/ferromagnetic structures[END_REF][START_REF] Lei | Electric field control of domain wall logic in piezoelec-tric/ferromagnetic nanodevices[END_REF][START_REF] More-Chevalier | Magnetoelectric coupling in Pb(Zr,Ti)O3-Galfenol thin film heterostructures[END_REF][START_REF] Srinivasan | Composite magnetoelectrics: materials, structures, and applications[END_REF][START_REF] Cibert | Magnetoelectric coupling in Tb0.3Dy0.7Fe2/Pt/PbZr0.56Ti0.44O3 thin films deposited on Pt/TiO2/SiO2/Si substrate[END_REF] Alternatively and with great design flexibility, the significant ME coefficients are obtained, much larger than in single-phase ME materials. In particular, electric field controlled CME coupling effect has been widely studied, and devices are prepared for power-efficient voltage writing of magnetic bits, voltage tunable inductors, tunable bandpass filters, tunable phase shifters, etc.Conventional CME composites are the combination of FM materials and FE materials, in which the magnetic properties of FM materials can be manipulated by applying the voltage to FE materials through strain transferring. Recently, novel multiferroics composites of FM/AFE materials have attracted the researchers' attention.

Figure 2 . 11

 211 Figure 2.11 Magnetization as functions of temperature for different values of the electric (E) fields.[START_REF] Sosa | A study of the dielectric and magnetic properties of multiferroic materials using the Monte Carlo method[END_REF]

  dependent magnetization dynamics model, based on simultaneous evaluation of the stochastic Landau-Lifshitz-Bloch equation and heat flow equation, together with stochastic thermal fields and magnetoelastic contributions, as given in Fig.2.12. With stress values ranging from 80 to 180 MPa and maximum temperatures normalized to the Curie temperature ranging from 0.65 to 0.99, the thermally activated strain-induced magnetization switching can be realized.

Figure 2 .

 2 Figure 2.12 Heat-assisted strain-induced magnetization switching. (a) Heating and cooling of multiferroic layers during and after a laser pulse. The inset shows the temperature profile in the free layer; (b,c) heat-assisted magnetization switching starting from states "1" and "0" with two different stress directions.[START_REF] Lepadatu | Heat-assisted multiferroic solid-state memory[END_REF]

Figure 2 . 13

 213 Figure 2.13 Induced magnetic moment variance with (a) the bias magnetic field, (b)the bias electric field, (c) the angle θ between the sample plane and magnetic field, and (d) the thickness ratio between Terfenol-D and PLZST at a resonance frequency of 365 kHz.[START_REF] Zhou | Magnetoelectric coupling in antiferroelectric and magnetic laminate composites[END_REF]

Figure 2 . 14

 214 Figure 2.14The magneto-optical Kerr effect signal converted to the oscillation amplitude of the magnetization vector (∆γ) measured at different AC voltages.[START_REF] Mirshekarloo | Magneto-optical Kerr effect investigation on magnetoelectric coupling in ferromagnetic/antiferroelectric multilayer thin film structures[END_REF]

Figure 2 .

 2 Figure 2.15 (a) M-H hysteresis loops of the austenite and martensite NMG thin films, (b) Out-of-plane P-E hysteresis loops of the PLZST films with austenite andmartensite NMG film on the top, before and after annealing, respectively.[START_REF] Mirshekarloo | Ferroelastic strain induced antiferroelectric-ferroelectric phase transformation in multilayer thin film structures[END_REF]

Figure 2 . 16

 216 Figure 2.16 Enlarged view of PLZT (001) peaks measured under various tilt angles for (a) the as-deposited, and (b) the annealed samples, respectively. (c) d-spacing of PLZT (001) plane as a function of sin2ψ for the two samples.[START_REF] Wang | Coupling of ferroelastic strain and ferroelectric phase transition in NiMnGa/Pb0.97La0.02(Zr0.95Ti0.05)O3 bilayered films[END_REF]

Figure 2 .

 2 Figure 2.17 (a) The schematic of FeGaB film deposited on the top or on the side of PSZT ceramics, (b) P-E loop of PZST ceramic, (c) FMR spectra under varying Efield of FeGaB(top)/PSZT composites and inset the FMR field dependence of E-field, and (d) magnetization switch of FeGaB(top)/PSZT at 40.5 Oe induced by voltageimpulse.[START_REF] Zhou | Strong non-volatile voltage control of magnetism in magnetic/antiferroelectric magnetoelectric heterostructures[END_REF]

Figure 3 . 1 AFigure 3 . 2

 3132 Figure 3.1 A principle schematic representation of RF magnetron sputtering[103]

  3.3. A voltage supply (High voltage supply, SRS Model PS350) was equipped to apply an electric field during the measurements. The sample is put into VSM while two Cu wires are attached to the electrodes of the sample.

Figure 3 . 3

 33 Figure 3.3 (a) Home-made VSM system, (b) Connected sample.

4. 1 . 1 Introduction

 11 In strain-mediated multiferroic composites, the FE materials are usually materials with considerable E-induced strain, such as Pb(Zr, Ti)O3 ceramics and Pb(Mg1/3Nb2/3)O3-PbTiO3 single crystal.[START_REF] Yang | Voltage tunable multiferroic phase shifter with YIG/PMN-PT heterostructure[END_REF][START_REF] Tatarenko | Microwave magnetoelectric effects in ferrite-piezoelectric composites and dual electric and magnetic field tunable filters[END_REF][START_REF] Pettiford | Bias field effects on microwave frequency behavior of PZT/YIG magnetoelectric bilayer[END_REF][START_REF] Thiele | Influence of strain on the magnetization and magnetoelectric effect in La0.7A0.3MnO3 ∕PMN -PT (001) (A =Sr, Ca)[END_REF] PZT ceramics are the most wellknown ones because of their excellent piezoelectric properties. PMN is a representative of relaxation ferroelectric material due to its high electrostriction, relaxation behavior, and high dielectric constant. Therefore, the solid solution between PMN and PZT has superior dielectric, piezoelectric, electromechanical coupling, and electrostrictive properties than single PMN and PZT.[START_REF] Koval | Effect of PMN modification on structure and electrical response of xPMN-(1-x)PZT ceramic systems[END_REF][START_REF] Yimnirun | Effects of uniaxial stress on dielectric properties lead magnesium niobate-lead zirconate titanate ceramics[END_REF][START_REF] Ouchi | Piezoelectric properties of Pb(Mg1/3Nb2/3)O3-PbTiO3-PbZrO3 solid solution ceramics[END_REF]

Figure 4 . 1 4 . 1 . 2

 41412 Figure 4.1 Piezoelectric coefficient d33, electromechanical coupling factor kp, and quality factor Qm of the x PMN-(1-x) PZT ceramics.[START_REF] Wang | Effect of PMN content on the phase structure and electrical properties of PMN-PZT ceramics[END_REF]

Figure 4 . 2 Figure 4 . 3

 4243 Figure 4.2 (a) AFM picture of the polished PMN-PZT substrate surface, (b) XRD patterns of PMN-PZT ceramic powders.

Figure 4 . 4

 44 Figure 4.4 Schematic of PMN-PZT ceramic substrates for electrical measurements.

  Figure 4.5 shows the P-E hysteresis loop of PMN-PZT ceramic substrate. The remnant polarization is 34 μC/cm 2 , and the saturate polarization is 38 μC/cm 2 , indicating strong ferroelectricity in PMN-PZT. The coercive electric field (EC) of PMN-PZT is low, about 4.2 kV/cm.

Figure 4 . 5

 45 Figure 4.5 The P-E loop of PMN-PZT ceramic substrate.

Figure 4 . 6

 46 Figure 4.6 (a) The in-plane strain and (b) out-of-plane strain of PMN-PZT ceramic substrate.

  [START_REF] Lu | Electric-field-induced AFE-FE transitions and associated strain/preferred orientation in antiferroelectric PLZST[END_REF] Such high strains have been considered to be related to a reversible, E-field-induced, tetragonal antiferroelectric (AFET) to rhombohedral ferroelectric (FER) phase transition.[START_REF] Blue | In situ x-ray diffraction study of the antiferroelectric-ferroelectric phase transition in PLSnZT[END_REF]

Figure 4 . 7

 47 Figure 4.7 Ternary phase diagram of the Pb0.97La0.02(Zr, Sn, Ti)O3 system.[121]

Figure 4 . 8

 48 Figure 4.8 (a) The AFM picture and (b) XRD pattern of PLZST ceramic substrate.

Figure 4 .

 4 9 illustrates the surface of the PLZST ceramic substrate with a scanning view of 200 μm. In this work, the relative density of PLZST ceramics is 96.8% through the Archimedes method, which is quite high when compared with other ceramics prepared by conventional solid-state reaction.[START_REF] Simon-Seveyrat | Reinvestigation of synthesis of BaTiO3 by conventional solid-state reaction and oxalate coprecipitation route for piezoelectric applications[END_REF][START_REF] Anselmi-Tamburini | Solid-state synthesis and spark plasma sintering of submicron BaYxZr1-xO3-x/2 (x=0, 0.08 and 0.16) ceramics[END_REF][START_REF] Galassi | Influence of processing parameters on the properties of PZT materials[END_REF]

Figure 4 . 9

 49 Figure 4.9 The surface SEM picture of the PLZST ceramic substrate.

Figure 4 .

 4 Figure 4.10 The P-E loop and S-E curve of the PLZST ceramic substrate.

Figure 5 .

 5 1 shows the XRD patterns of NMG film annealed at different temperatures. It can be clearly observed that the austenitic NMG film is crystallized at 500 ℃ where the sharp Bragg peak (2θ ~ 43°) presents. The low FWHM of NMG film also indicates the good crystallization happens at 500 ℃.

Figure 5 . 1

 51 Figure 5.1 The XRD patterns of NMG films on Si substrates (A: Austenite; M: Martensite).

Figure 5 . 2

 52 Figure 5.2 The magnetic hysteresis loops of NMG films on Si substrates annealed at different temperatures.

Figure 5 . 3 Figure 5 . 4

 5354 Figure 5.3 Schematic of NMG/PLZST/NMG multiferroic heterostructure.

Figure 5 . 5

 55 Figure 5.5 The M-H loop of NMG/PLZST/NMG heterostructure.

Figure 5 . 6

 56 Figure 5.6 The M-H loop of NMG/PLZST/NMG heterostructure with and without an electric field.

Figure 5 . 7

 57 Figure 5.7 (a) (ΔM/M0)-E curves under bias magnetic field of 0, 100 and 500 Oe, and (b) S-E curve and |ΔM/M0|-E loop measured at 500 Oe of NMG/PLZST/NMG heterostructure.

Figure 5 . 8

 58 Figure 5.8 The (ΔMmax/M0)-H curve of NMG/PLZST/NMG heterostructure.

Figure 5 .

 5 Figure 5.8 shows the curve of E-induced ΔM/M0 changing with the bias magnetic field. It can be seen that the maximum of ΔM/M0 (15%) is reached at 0 Oe, which is a very important asset in terms of design simplification and cost reduction, thus suggesting significant potential for future ME devices. With the bias magnetic field increases, the ΔM/M0 reduces quickly, indicating that it becomes difficult for the electric field induced strain to destroy the alignment of well-ordered magnetic domains and rotate them to other directions.

Figure 5 . 9

 59 Figure 5.9 (a) ΔM/M0 and CME coefficient versus E under a bias magnetic field of 500 Oe, (b) the maximum of CME coefficient versus bias H under positive and negative voltages of NMG/PLZST/NMG sample.

Figure 6 . 1

 61 Figure 6.1 (a) The XRD patterns and (b) M-H loops of YIG films on SiO2/Si substrates with different annealing temperatures.

Figure 6 . 2

 62 Figure 6.2 Schematic of YIG/Pt/PLZST (PMN-PZT)/Pt heterostructure (a) front side and (b) backside.

Figure 6 . 3

 63 Figure 6.3 Schematic of (a) YIG/Pt/PLZST/Pt and (b) YIG/Pt/PMN-PZT/Pt ME heterostructures.

Figure 6 .

 6 3 illustrates the schematic of YIG/Pt/PLZST/Pt and YIG/Pt/PMN-PZT/Pt with the application of the electric and inplane magnetic field.

Figure 6 . 4 6 . 2 YIG/Pt/PLZST/Pt heterostructure 6 . 2 . 1

 6462621 Figure 6.4 The XRD pattern and inset cross-sectional SEM image of (a) YIG/Pt/PLZST/Pt and (b) YIG/Pt/PMN-PZT/Pt heterostructures.

Figure 6 . 5

 65 Figure 6.5 The in-plane M-H loops of YIG/Pt/PLZST/Pt heterostructure.

Figure 6 . 6

 66 Figure 6.6 (a) M-H loops, inset the enlarged M-H loops, (b) HC-E curve of YIG/Pt/PLZST/Pt heterostructure.

Figure 6 . 7

 67 Figure 6.7 (a) χ-H curves with different electric fields and (b) χ-E curve at 10 Oe of YIG/Pt/PLZST/Pt heterostructure.

  excluded, and the YIG/Pt/PLZST/Pt heterostructure was fully magnetized at the same high bias magnetic field (2 T) to ensure the same initial state. The M-E curves under the bias magnetic field of 0, 15, and 30 Oe of YIG/Pt/PLZST/Pt heterostructure are given in Fig.5.13. Under a specific magnetic field, the cycle for the change of the electric field is as follows: 0 → 54 → -54 → 0 kV/cm. The M-E curves were measured in two electrical cycles with a bias magnetic field of 0 Oe, while only one cycle for 15 and 30 Oe.

Figure 6 . 8

 68 Figure 6.8 The M-E curves of YIG/Pt/PLZST/Pt at 0 Oe, 15 Oe, and 30 Oe. The labeled M0i and M0f are the initial magnetization and final magnetization at 0 kV/cm.

Figure 6 .

 6 Figure 6.8 (a) presents two distinct effects. One is the reversible strain-driven magnetization change. The horn-like M-E curves show a similar shape to the S-E curve of PLZST in Fig.4.10, demonstrating the magnetization variations are strongly induced electro-strain originated from PLZST ceramic. A second effect appears which also induces a change in the magnetization, that is, the difference of the magnetization between the initial state (M0i), at the beginning of the measurements, and the magnetization state at the end of the first electrical cycle (M0f). Since the YIG/Pt/PLZST/Pt heterostructure was fully magnetized at 2 T before the measurements, the magnetic domains of YIG film are highly aligned to one direction. The value of M0i

Figure 6 . 9

 69 Figure 6.9 (a) The relative magnetization change versus magnetic field (∆M/M0f -H) curve, (b) (M0i-M0f)/M0f -H curve of YIG/Pt/PLZST/Pt heterostructure.

Figure 6 . 10

 610 Figure 6.10 The magnetization versus electric field curves of YIG/Pt/PLZST/Pt (a) 0 Oe, (b) 30 Oe.

Figure 6 .

 6 Figure 6.11 ∆M/M0-E and αCME-E curves at 0 Oe of YIG/Pt/PLZST/Pt heterostructure.

Figure 6 .

 6 Figure 6.12 The CME coefficient versus bias magnetic field curves of YIG/Pt/PLZST/Pt heterostructure.

Figure 6 . 13

 613 Figure 6.13 The out-of-plane M-H loops of YIG/Pt/PLZST/Pt heterostructure with different electric fields.

Figure 6 . 14

 614 Figure 6.14 The out-of-plane M-E curves of YIG/Pt/PLZST/Pt heterostructure (a) 500 Oe, (b) 1000 Oe.

Figure 6 .

 6 Figure 6.15 (a) The in-plane M-E curve and (b) out-of-plane M-E curve of YIG/Pt/PLZST/Pt heterostructure.

Figure 6 . 16

 616 Figure 6.16 Schematic of the veer of the anisotropy axis in YIG film.

Figure 6 . 17 Table 6 . 5

 61765 Figure 6.17 The normal and inset the enlarged M-H loop of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 .

 6 Figure 6.18 (a) M-H loops with different electric fields, inset the enlarged M-H loops, (b) Mr-E and HC-E curves of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 .

 6 19 (a) presents the χ-H curves under electric fields of -4, 0 and 12 kV/cm in YIG/Pt/PMN-PZT/Pt heterostructure. Under a specific electric field, the χ reaches the maximum at HC. The 72 HC varies with electric fields, thus the E-field controlled χ peaks at different magnetic fields (20 ~30 Oe). Under a specific magnetic field of 15, 20 and 25 Oe, the E-field controlled χ variations are given in Fig.6.6 (b). At 20 Oe, it can be seen that the χ changes dramatically with the external electric field, being 33.6 at 12 kV/cm, 27.6 at zero fields and 15.5 at the coercive electric field. This equals a large tunable susceptibility ∆χ = 18.1 and ∆χ/χ0 =65.6% under a low magnetic field of 20 Oe, implying great potential in ME tunable devices. When the magnetic field is 15 Oe or 25 Oe, the tunable susceptibility ∆χ/χ0 can reach 80.4% and 28.5%, respectively.

Figure 6 .

 6 Figure 6.19 (a) χ-H curves with different electric fields and (b) χ-E curve at 15, 20, and 25 Oe of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 .

 6 [START_REF] Liu | Voltage tuning of ferromagnetic resonance with bistable magnetization switching in energyefficient magnetoelectric composites[END_REF] shows the first and second cycles of M-E curves with a bias magnetic field of 0 Oe as a representative. The difference of the magnetization between M0i and M0f exists in the first cycle of the M-E curve, which is due to the irreversible demagnetization of the previous "poling process" of YIG film.The first cycle of the M-E curve is measured from a high positive electric field of +12 kV/cm, so the curve is affected by this measurement process and presents in an 73 asymmetric butterfly-like shape. After the first cycle of M-E measurement, the electrical and magnetic load history can be removed. Then the second cycle of M-E measurement can be carried out to fully reveal the CME coupling effect in YIG/Pt/PMN-PZT.

Figure 6 . 20

 620 Figure 6.20 The first and second cycle of M-E curves at 0 Oe in YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 21

 621 Figure 6.21 The electric-field-dependent magnetization (M-E) in YIG/Pt/PMN-PZT/Pt under the bias magnetic field of 0, 10, 25, 50, 75, 100, 200 and 500 Oe

Figure 6 .

 6 Figure 6.22 (a) ∆M/M0-E and αCME-E curves at 0 Oe, and (b) |∆M/M0| -H and (αCME) -H curves of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 23 Figure 6 .

 6236 Figure 6.23 The out-of-plane strain evolution of PMN-PZT ceramic substrate.

Figure 6 .

 6 Figure 6.25 M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure under cycling electric field with different amplitudes at 0 Oe. (a) Poled at +12 kV/cm and (b) poled at -12 kV/cm.

Figure 6 .

 6 Figure 6.26 χ-E curves of YIG/Pt/PMN-PZT/Pt heterostructure under cycling electric field with different amplitudes at 20 Oe.

Figure 6 .

 6 Figure 6.27 (a) The M-E curve with an electric field cycle from -4.6 kV/cm to 4.6 kV/cm, (b) the time-dependent magnetization switched by the electric field at 0 Oe of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 28

 628 Figure 6.28 The χ-E curve with an electric field cycle from -4 kV/cm to 4 kV/cm at 20 Oe of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 29

 629 Figure 6.29 The electric-field dependent M-H loops of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 30

 630 Figure 6.30 The out-of-plane M-E curves of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 31

 631 Figure 6.31 The in-plane M-E curve at 0 Oe and out-of-plane M-E curve at 500 Oe of YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 6 . 32

 632 Figure 6.32 Schematic of electric field induced anisotropy change in YIG/Pt/PMN-PZT/Pt heterostructure.

Figure 7 . 1

 71 Figure 7.1 (a) Schematic of TCFC multilayers and (b) TCFC/PLZST/Au heterostructure.

Figure 7 . 2

 72 Figure 7.2 (a) Front side of TCFC/PLZST (PMN-PZT)/Au heterostructure, (b) picture of TCFC/PLZST (PMN-PZT)/Au connected to sample holder, (c) front side and (d) backside of TCFC/PLZST (PMN-PZT)/Au connected with wires.

Figure 7 . 3

 73 Figure 7.3 The M-H loops of TCFC/PLZST/Au heterostructure along with different directions.

Figure 7 .

 7 Figure 7.4 shows the M-H loop of TCFC on PLZST ceramic along the hard axis.

Figure 7 . 4

 74 Figure 7.4 The M-H loop of TCFC/PLZST/Au heterostructure along the hard axis.

  ceramic substrate. The horn-like M-E curves demonstrate the effectiveness of the strain-mediated ME coupling effect. The application of strain (AFE-FE) can destroy

3 .

 3 Since the second cycles of M-E curves give the reversible E-field controlled magnetization switching, the ∆M and CME coefficient is calculated from the second half M-E curves, as given in Fig.7.6. The M0 is the magnetization without the electric field in the second half M-E curves, and the ∆M is the magnetization difference between ME and M0. The peaks of the CME coefficient always occur at the switching field of PLZST ceramic, where the phase transition strain bursts out. The quick change of Efield induced strain results in sharp magnetization change and leads to a large αCME at switching field of PLZST.

Figure 7 . 5

 75 Figure 7.5 The M-E curves of TCFC/PLZST/Au heterostructure along the hard axis at 0, 50, 100, 190, 500, 1000, and 1500 Oe.

Figure 7 . 6

 76 Figure 7.6 The ∆M-E, and αCME-E curves of TCFC/PLZST/Au heterostructure along the hard axis at 0, 50, 100, 190, 500, 1000, and 1500 Oe.

Figure 7 . 7

 77 Figure 7.7 (a) The ∆M/M0-H curve and (b) αCME-H curve of TCFC/PLZST/Au heterostructure along hard axis.

Figure 7 .

 7 Figure 7.7 shows the ∆M/M0-H curve and αCME-H curve of TCFC/PLZST/Au heterostructure along the hard axis. The relative magnetization change ∆M/M0 can reach -9.8% with an absence of a magnetic field, then it goes up to 1.9% at 500 Oe and

Figure 7 .

 7 Figure 7.8 shows the M-H loops along the easy axis of TCFC/PLZST/Au heterostructure under different electric fields. With an application of the electric field of 44 kV/cm, the electric-field induced AFE-FE phase transition happens in PLZST, accompanied by a large phase deformation strain. However, the large E-induced strain did not lead to a distinct change in M-H loops. The inset shows enlarged M-H loops around HC, but the E-induced variation is too small to be distinguished. The highly intrinsic crystallographic orientation in the magnetic field of TCFC makes it difficult to induce the magnetic moment rotation or motion by strain. The strain-induced magnetization variations are almost forbidden along the easy axis. Thus it is difficult to realize the CME effect along the easy axis in TCFC/PMN-PZT/Au heterostructure.

Figure 7 . 8

 78 Figure 7.8 The M -H loops along the easy axis of TCFC/PLZST/Au heterostructure under different electric fields.

Figure 7 .

 7 Figure 7.9 shows the M-H loops along the OOP direction of TCFC/PLZST/Au heterostructure under different electric fields. With an application of the electric field, the M-H loops have not presented any visible variations. The TCFC film has a strong

Figure 7 . 9

 79 Figure 7.9 The M -H loops along the OOP direction of TCFC/PLZST/Au heterostructure under different electric fields.

Figure 7 .

 7 Figure 7.10 The M-H loops of TCFC/PMN-PZT/Au heterostructure along with different directions.

Figure 7 .

 7 Figure 7.11 shows the M-H loops along the hard axis of TCFC/PMN-PZT/Au heterostructure under different electric fields. The shape of M-H loops changes with different electric fields, indicating apparent strain-induced magnetization variations

Figure 7 .

 7 Figure 7.11 The M-H loops along the hard axis of TCFC/PMN-PZT/Au heterostructure under different electric fields.

Figure 7 . 12 101 Figure 7 . 13

 712101713 Figure 7.12 The M-E curves along the hard axis of TCFC/PMN-PZT/Au heterostructure at 0 and 40 Oe.

Figure 7 .

 7 Figure 7.13 shows M-E curves along the hard axis of TCFC/PMN-PZT/Au heterostructure under different magnetic fields. With the absence of an external magnetic field, the M-E curve has a typical butterfly shape which resembles the S-E curves of PMN-PZT ceramic. With a magnetic field of 40 Oe, the shape of the M-E curve begins to change, and the peak of hindwing of the butterfly-like M-E curve gets inverted. Then as the magnetic field increases, the inversed part expands. At 500 Oe, the whole hindwing of the M-E curve is inverted when compared with the butterfly-like

Figure 7 . 14

 714 Figure 7.14 The ideal M-H loops along the hard axis of a single crystal with different strains.

Figure 7 .

 7 Figure 7.15 The M-H loops along the hard axis of TCFC/PMN-PZT/Au heterostructure under different magnetic fields.

  7.16. The peaks of the CME coefficient always occur around the EC of PMN-PZT ceramic, where the large strain presents. The quick change of E-field induced strain results in sharp magnetization change and leads to a large αCME around the EC of PMN-PZT.

Figure 7 .

 7 Figure 7.16 The ∆M-E and αCME-E curves of TCFC/PMN-PZT/Au heterostructure along the hard axis at 0, 40, 80, 160, 300, 500, 800, 1000, 1300, 1500 and 2800 Oe.

Figure 7 .

 7 Figure 7.17 The 𝛼 𝑚𝑎𝑥 -H curve along the hard axis of TCFC/PMN-PZT/Au heterostructure.

Figure 7 .Figure 7 . 18

 7718 Figure 7.[START_REF] Vopson | Fundamentals of Multiferroic Materials and Their Possible Applications[END_REF] shows the M-H loops along the easy axis of TCFC/PMN-PZT/Au heterostructure under different electric fields. With an application of the electric field of 4 kV/cm, the maximum in-plane strain can be generated in PMN-PZT ceramic, but the large tensile strain did not lead to a distinct change in M-H loops. The inset shows minor variations in HC, but the variation is too small. The highly intrinsic crystallographic orientation in the magnetic field of TCFC makes it difficult to induce the magnetic moment rotation or motion by strain. The strain-induced magnetization variations are almost forbidden along the easy axis. Thus it is difficult to realize the CME effect along the easy axis in TCFC/PMN-PZT/Au heterostructure.

Figure 7 .

 7 Figure 7.19 shows the M-H loops along the OOP direction of TCFC/PMN-PZT/Au heterostructure under different electric fields. With an application of different electric fields, the linear-like M-H loops vary with the E-induced strain. To reveal the relationship between the slope of M-H loops and the electric field, a linear fitting of M-H loops has been done, as given in Fig. 7.20 (a). Then the fitting slope of M-H loops versus the electric field curve can be carried out in Fig. 7.20 (b).

Figure 7 .

 7 Figure 7.19 The M-H loops along the OOP direction of TCFC/PMN-PZT/Au heterostructure under different electric fields.

Figure 7 .

 7 Figure 7.20 (a) The M-H loops and linear fitted under different electric fields, and (b) the slope of fitted lines versus electric fields along the OOP direction of TCFC/PMN-PZT/Au heterostructure.

  Figure 7.21 shows the M-E curves of TCFC/PMN-PZT/Au heterostructure at 5000 and 15000 Oe. The typical butterfly-like M-E curves can be well observed.

Figure 7 .

 7 Figure 7.21 The M-E curves along the OOP direction of TCFC/PMN-PZT/Au heterostructure (a) at 5000 Oe and (b) at 15000 Oe.

Figure 7 .

 7 Figure 7.22 The ∆M/M0-E and αCME-E curves along the OOP direction of TCFC/PMN-PZT/Au heterostructure (a) at 5000 Oe and (b) at 15000 Oe.

  [(TbCo2)/(FeCo)]20/PLZST/Au and [(TbCo2)/(FeCo)]20/PMN-PZT/Au heterostructures have been fabricated. To explore the CME coupling effect along different orientations, the TCFC films are prepared to exhibit mutually perpendicular in-plane easy axis and hard axis. The E-induced magnetism along easy axis, hard axis, and OOP direction have been discussed.

Figure 8 . 1 8 . 1 . 2

 81812 Figure 8.1 The XRD patterns of LNO films on SiO2/Si substrates.

Figure 8 . 2

 82 Figure 8.2 The XRD patterns of PZ films on LNO/SiO2/Si substrates.

Figure 8 . 3

 83 Figure 8.3 Schematic of Pt/PZ/LNO/SiO2/Si heterostructure.

  Figure 8.4 (b) illustrates four dielectric constant (εr) peaks, as well as four tanδ peaks, which are the typical characteristics of AFE materials.

Figure 8 . 4

 84 Figure 8.4 (a) The P-E loops and (b) εr-E and tanδ-E curves of Pt/PZ/LNO/SiO2/Si heterostructure.

Figure 8 . 5

 85 Figure 8.5 The XRD patterns of PZT films on LNO/SiO2/Si substrates.

Figure 8 . 6

 86 Figure 8.6 Schematic of Pt/PZT/LNO/SiO2/Si heterostructure.

Figure 8 .

 8 Figure 8.7 shows the electrical properties of Pt/PZT/LNO/SiO2/Si heterostructure. The typical ferroelectric hysteresis P-E loop is well observed in Fig. 8.7 (a), indicating the ferroelectricity of PZT films. The εr-E and tanδ-E curves of Pt/PZT/LNO/SiO2/Si heterostructure are also demonstrated the ferroelectric properties of PZT films.

Figure 8 . 7 8 . 1 . 4

 87814 Figure 8.7 (a) The P-E loops and (b) εr-E and tanδ-E curves of Pt/PZT/LNO/SiO2/Si heterostructure.

Figure 8 . 8

 88 Figure 8.8 Schematic of Pt/PZ/LNO/SiO2/Si heterostructure.

Figure 8 . 9 ( 1 )

 891 Figure 8.9 XRD patterns of (a) YIG/Pt/PZ/LNO/SiO2/Si and (b) YIG/Pt/PZT/LNO/SiO2/Si heterostructures.

Figure 8 .Table 8 . 5

 885 Figure 8.10 In-plane M-H loops of YIG/SiO2/Si and YIG/Pt/PZ/LNO/SiO2/Si heterostructures.

Figure 8 .

 8 Figure 8.11 (a) The MFM system, and (b) schematic of the YIG/Pt/PZ/LNO/SiO2/Si heterostructure with probes.

Figure 8 .

 8 Figure 8.12 (a) AFM surface morphology and (b) MFM image of YIG films on SiO2/Si substrate.

(

  33%) with a low magnetic field of 10 Oe in YIG/Pt/PLZST/Pt heterostructure. The in-plane CME coefficient in YIG/Pt/PMN-PZT/Pt heterostructure can reach 16.95 × 10 -8 s/m in the absence of bias magnetic field at EC of PMN-PZT ceramic, and the maximum of in-plane CME coefficient (18.15 × 10 -8 s/m) is observed under the magnetic field of 25 Oe. When the application of electric field amplitude is close EC of PMN-PZT, the voltage-impulse reversible magnetization switching and the linear E-tunable susceptibility change can be realized in YIG/Pt/PMN-PZT/Pt heterostructures. Forth, the CME coupling effect of the rare-earth Tb modified 3d-transition alloys (TbCo2)/(FeCo) coupled with PLZST and PMN-PZT ceramic substrates has also been revealed. The [(TbCo2)/(FeCo)]20/PLZST/Au and [(TbCo2)/(FeCo)]20/PMN-PZT/Au heterostructures have been fabricated, and the CME coupling effect along different directions have been discussed. The E-induced magnetization variation is suppressedalong the easy axis of TCFC film by the highly intrinsic crystallographic orientation with the magnetic field. Along the hard axis, a reverse process of M-E curves appears with the bias magnetic field increasing. The appearances of the inverse transition of M-E curves may be due to the uneven distribution of TCFC film on a ceramic substrate with inevitable holes. In TCFC/PLZST/Au, the maximum of ∆M/M0 can reach -9.8%

3 .

 3 In YIG/AFE and YIG/FE heterostructures, a follow-up study should investigate the influence of the thickness of the Pt electrode in the CME coupling effect. We have only shown the CME coupling effect in YIG/Pt/AFE/Pt and YIG/Pt/FE/Pt heterostructures, where the thickness of the Pt layer is 80 nm. A more detailed study of the interlayer thickness in the CME coupling effects and more widely, the interface effects should be carried out. 4. A subsequent study involving the fabrication of FM-AFE and FM-FE multiferroic films should be conducted and compared. The AFE-and FE-based multiferroic heterostructures in a thin-film form with differently-oriented AFE and FE films should be investigated, and therefore the CME coupling effect can be explored, which is showing potential in a variety of practical applications.In closing, the current work suggests a variety of research scenarios-both in AFEand FE-based multiferroic heterostructures and in the characterization of E-field controlled magnetic properties at static state. Based on the specific characteristics of each component, we can study the electrical, magnetic, and magnetoelectric properties of the self-assembled heterostructures, and more importantly, the comparison of AFE-

  large strain change from E-induced AFE-FE phase transition in AFE materials offers themselves great advantages to play a key role in realizing voltage control of magnetism. The present research of novel FM/AFE multiferroic structures has revealed a new path of design CME composites.

Since hundreds of literature focused on FE-based ME composites have been published, and various heterostructures have been brought out, the current research on AFE-based ME composites is quite limited. Only a few AFE-based ME heterostructures have been developed, and the published papers are less than ten. Thus, more AFE-based ME composites need to be developed. Besides, the difference of AFEand FE-based ME composites in CME coupling effect is still unclear, which is necessary to understand these two type of heterostructures. Many applications of FM/FE composites have already been proposed, while few are related to AFE-based ME composites. The exploration of potential applications in AFE-based ME composites is also needed to be done.

Table 3 . 1

 31 Main raw materials for the preparation of ceramics

	Samples	Raw powders	Purity
		Pb3O4	99.26%
	0.25(PbMn1/3Nb2/3O3)-	MgCO3	99.98%
	0.75(PbZr0.47Ti0.53O3)	Nb2O5	99.81%
	(PMN-PZT)	TiO2	99.38%
		ZrO2	99.99%
		Pb3O4	99.26%
		La2O3	99.95%
	(Pb0.97La0.02)(Zr0.6Sn0.3Ti0.1)O3 (PLZST)	ZrO2	99.99%
		SnO2	99.5%
		TiO2	99.38%

CME Coupling Effect in NMG/AFE Heterostructure 5.1 Introduction

  good electric field-induced strain in PMN-PZT and PLZST ceramics gives us a way to fabricate FE-and AFE-based ME composites based on these ceramic substrates and reach good CME coupling effect. Moreover, the comparison of the performance of FE-and AFE-based ME composites can be carried out subsequently.

	5 Previous researchers have studied a variety of layered ME heterostructures of FM
	and FE components.[7,20-22] Actually, the main interest has been focused on FM/FE
	heterostructures and in particular the electric-field tuning of magnetic response
	resulting from the strain of FE component transferred to FM component, such as
	CoFe2O4/BaTiO3, Ni0.8Zn0.2Fe2O4/Pb(Zr,Ti)O3, Terfenol-D/Pb(Mn,Nb)O3-PbTiO3,
	FeBSiC/PVDF and so on. [16-19] AFE materials, with novel properties of electric-

field-induced AFE/FE phase transition, exhibit a large and sharp strain change during

Table 5 . 1

 51 Sputtering parameters of NMG films

	Sputtering Condition	NMG
	Vacuum (Pa)	< 3 ×10 -4
	Power (W)	60
	Temperature (℃)	RT
	Atmosphere	Ar
	Pressure (Pa)	1

Table 5 . 2

 52 Annealing parameters of NMG films

	Annealing Condition	NMG

CME Coupling Effect of YIG Films on AFE and FE Ceramics Ferrite

  Even though, these CME results also reveal the great potential of AFE materials in the ME coupling effect. It is the first time that the value of the CME coefficient is presented for an FM/AFE heterostructure.6 CME Coupling Effect of YIG Films on AFE and FE Ceramics materials are an essential family of magnetic materials, and they are also widely used in ME coupling composites. Among them, YIG is considered as a good choice for the magnetic material when targeting tunable multiferroic microwave devices thanks to its outstanding advantages including small magnetocrystalline anisotropy, narrow ferromagnetic resonance linewidth, high resistivity, and low microwave losses. Therefore, many works are focused on magnetoelectric interactions of YIG/FE. Since YIG has preponderant advantages in ME composites, but the most work of YIG/FE composites is focused on the E-field controlled FMR shift, which is an indirect method to characterize the CME coupling effect. The most direct

55

6

manifestation of the CME effect, the E-field controlled magnetization variations, is less studied. To understand the converse magnetoelectric (CME) coupling effect, it is necessary to study the electric-field tuning of magnetism of YIG at a static state.

[START_REF] Lian | Electric field tuning of magnetism in heterostructure of yttrium iron garnet film/lead magnesium niobate-lead zirconate titanate ceramic[END_REF]

Table 6 . 1

 61 Sputtering parameters of YIG films

	Sputtering Condition	YIG

Table 6 . 2

 62 Annealing parameters of YIG films

	Annealing Condition	YIG
	Heating Rate (℃/min)	2
	Annealing (℃)	700, 750
	Duration (min)	60

Table 6 .

 6

	Ms	Mr	HC	χ
	(emu/cm 3 )	(emu/cm 3 )	(Oe)
	53.8	37.3	15	41

3

The values of Ms, Mr, HC, and χ of YIG/Pt/PLZST/Pt heterostructure.

Table 6 . 4

 64 Summary of magnetoelectric coupling parameters for some laminated structures.

	Structure	αCME (×10 -8 s/m)	∆M/M	bias H (Oe)	∆χ/χ0
	YIG/PLZST (this work)	11.6	17.4%	0	32.7% @10 Oe
	YIG/PMN-PZT[130]	10.9	7%	40	16.8% @30 Oe
	LSMO/PMN-PT(001)[111]	6	25%	100	-
	CFO/PMN-PT(001)[131]	3.2	6%	500	-
	CFB/ PMN-PT(001)[132]	200	17.6%	5	-

For comparison purposes, Table

6

.4 summarizes important results measured at a static magnetic state for different layered structures recently reported. In review, we observe that the AFE-based structure reported here demonstrates the first reported CME coefficient and relative susceptibility change for ME heterostructures. Further, this

Table 7 .

 7 1 gives the reversible and irreversible parts of E-field induced magnetization in TCFC/PLZST/Au heterostructure along the hard axis. The irreversible part (M0i-M0f) is calculated from the first loop of M-E curves. With the external magnetic field increasing, the ∆M is firstly negative; then it goes to positive. The irreversible part is positive under a different magnetic field, demonstrating the existence of an unrecoverable part of the magnetic domain motion caused by E-field induced strain.

Table 7 . 1

 71 The reversible and irreversible parts of E-field induced magnetization in TCFC/PLZST/Au heterostructure along the hard axis.

	Bias H (Oe)	0	50	100	190	500	1000	1500
	∆M (emu/cm	3 )	-10.88 -10.33 -10.14	-7.45	6.35	12.28	7.58
	(M0i-M0f) (emu/cm	3 ) 7.51	7.52	7.70	10.27	14.04	6.46	1.36

Table 8 . 1

 81 Sputtering parameters of LNO films

	Sputtering Condition	LNO
	Vacuum (Pa)	< 3 ×10 -4
	Power (W)	100
	Temperature (℃)	450
	Atmosphere	O2/Ar=1/4
	Pressure (Pa)	1

Table 8 . 2

 82 Annealing parameters of LNO films

	Annealing Condition	LNO
	Heating Rate (℃/min)	2
	Annealing (℃)	750
	Duration (min)	60

Table 8 . 3

 83 Sputtering parameters of PZT films

	Sputtering Condition	PZT
	Vacuum (Pa)	< 3 ×10 -4
	Power (W)	100
	Temperature (℃)	RT
	Atmosphere	Ar
	Pressure (Pa)	1

Table 8 . 4

 84 Annealing parameters of PZT films

	Annealing Condition	PZT
	Heating Rate (℃/min)	2
	Annealing (℃)	750
	Duration (min)	30
	Cooling Rate (℃/min)	3
	Atmosphere	Air

  and FE-based CME coupling effects are given, which may hold promise for different multifunctional applications. Duan, J.P. Velev, R.F. Sabirianov, Z. Zhu, J. Chu, S.S. Jaswal, E.Y.

	ground	states,	Advanced	Materials.	21	(2009)	3470-3474.
	doi:10.1002/adma.200900278.
	[72] C.-G.
	Energy	Density	Capacitors,	Materials.	8	(2015)	8009-8031.
	doi:10.3390/ma8125439.

Tsymbal, Surface magnetoelectric effect in ferromagnetic metal films, Physical Review Letters. 101 (2008) 137201. doi:10.1103/PhysRevLett.101.137201.

CME Coupling Effect in NMG/AFE Heterostructure

kV/cm

kV/cm 4.6 kV/cm 6 kV/cm 12 kV/cm

) H (Oe)

YIG/AFE and YIG/FE Heterostructures in Full-Thin-Film Form

Acknowledgements

VII

Acknowledgements

At the end of my work, I'd like to take this opportunity and present my gratitude to a number of people in my life. It is their trust and support that made the presented work possible. First and foremost I would like to express my sincere appreciation to Prof. Pascal AUBERT and Prof. Gilles POULLAIN. Thanks for their patient and careful review, and their valuable comments are very important. I would also thank Madame Nathalie LEMEE for her helpful suggestion to improve our work. Her experiences in material fabrication has offered us many impressive thinking. Then I would like to express my heartfelt gratitude to my supervisor Prof. Denis Acknowledgements Acknowledgements IX Table of Contents X Acknowledgements .................................................................................................... VII Table of Contents ........................................................................................................ X