Felipe Luiz Alvares Vital
Growth of Fe/BaTiO
3heterostructures for in-situ
characterization of the magnetoelectric coupling
Crescimento de heteroestruturas de Fe/BaTiO
3para
caracteriza¸
c˜
ao in-situ do acoplamento
magnetoel´
etrico
CAMPINAS 2019
Growth of Fe/BaTiO
3heterostructures for in-situ
characterization of the magnetoelectric coupling
Crescimento de heteroestruturas de Fe/BaTiO
3para
caracteriza¸
c˜
ao in-situ do acoplamento
magnetoel´
etrico
Thesis presented to the Institute of Physics Gleb Wataghin of the University of Camp-inas in partial fulfillment of the requirements for the degree of Master in Physics.
Disserta¸c˜ao apresentada ao Instituto de F´ısica Gleb Wataghin da Universidade Es-tadual de Campinas como parte dos requi-sitos exigidos para o t´ıtulo de Mestre em F´ısica.
Supervisor/Orientador: Julio Criginski Cezar
Co-supervisor/Co-orientador: Marcos Cesar de Oliveira
Este exemplar corresponde `a vers˜ao final da tese defendida pelo aluno Felipe Luiz Alvares Vital, e orientada pelo Prof. Dr. Julio Criginski Cezar.
CAMPINAS 2019
Vital, Felipe Luiz Alvares,
V83g VitGrowth of Fe/BaTiO3 heterostructures for in-situ characterization of magnetoelectric coupling / Felipe Luiz Alvares Vital. – Campinas, SP : [s.n.], 2019.
VitOrientador: Julio Criginski Cezar.
VitCoorientador: Marcos César de Oliveira.
VitDissertação (mestrado) – Universidade Estadual de Campinas, Instituto de Física Gleb Wataghin.
Vit1. Acoplamento magnetoelétrico. 2. Multiferróico. 3. Microscopia eletrônica de fotoemissão. 4. Filmes finos. I. Cezar, Julio Criginski, 1972-. II. Oliveira, Marcos César de, 1969-. III. Universidade Estadual de Campinas. Instituto de Física Gleb Wataghin. IV. Título.
Informações para Biblioteca Digital
Título em outro idioma: Crescimento de heteroestruturas de Fe/BaTiO3 para
caracterização in-situ do acoplamento magnetoelétrico
Palavras-chave em inglês:
Magnetoelectric coupling Multiferroic
Photoemission electron microscopy Thin films
Área de concentração: Física Titulação: Mestre em Física Banca examinadora:
Julio Criginski Cezar [Orientador] Flávio Garcia
Varlei Rodrigues
Data de defesa: 30-09-2019
Programa de Pós-Graduação: Física
Identificação e informações acadêmicas do(a) aluno(a) - ORCID do autor: https://orcid.org/0000-0002-2056-6496 - Currículo Lattes do autor: http://lattes.cnpq.br/2118369148369472
INSTITUTO DE FÍSICA “GLEB WATAGHIN”, DA UNIVERSIDADE ESTADUAL DE CAMPINAS, EM 30 / 09 / 2019.
COMISSÃO JULGADORA:
- Prof. Dr. Julio Criginski Cézar – Orientador – LNLS
- Prof. Dr. Flávio Garcia – CBPF
- Prof. Dr. Varlei Rodrigues – DFA/IFGW/UNICAMP
OBS.: Ata da defesa com as respectivas assinaturas dos membros encontra-se no
SIGA/Sistema de Fluxo de Dissertação/Tese e na Secretaria do Programa da
Unidade.
CAMPINAS
2019
so much pain and here you are making gold out of it
Realmente foi um trabalho muito ´arduo obter esta disserta¸c˜ao de mestrado e n˜ao posso deixar de agradecer primeiramente meu orientador Julio Criginski Cezar por todo o apoio emocional e cient´ıfico que ele tem me proporcionado desde a minha primeira inicia¸c˜ao cient´ıfica que comecei em 2013. Sempre penso sobre o conceito de “pai cient´ıfico”, algu´em que te guie na carreira cient´ıfica como um pai mesmo. Eu sinto que eu n˜ao poderia ter escolhido um melhor para mim.
Agrade¸co tamb´em a meus pais, Rivkah Val´eria Alvares e Jorge Luiz Frederich Vital, por terem me dado a vida. Sempre estiveram dispon´ıveis pra me dar suporte emocional e financeiro. Principalmente `a minha m˜ae que me criou a vida inteira proporcionando todos os elementos necess´arios para que eu me tornasse um ser humano completo.
Fico muito agradecido a todo mundo do grupo SIMM e da PGM. Estamos em um ambi-ente de trabalho maravilhoso proporcionado pela constante presen¸ca de bolos e conversas boas. Agradecimentos especiais a Thiago Mori, Pedro Schio, Flavia Estrada, Dayane Chaves e Jeovani Brand˜ao, que fizeram contribui¸c˜oes important´ıssimas na obten¸c˜ao e discuss˜ao dos resultados deste trabalho.
Al´em disso, n˜ao posso deixar de mencionar a ala jovem do nosso grupo: Pedro Cae-tano, Paloma Jackson, Ma´ıra Neme, Lucas Godinho, Karine Alcˆantara e Lidiana Moraes. Sempre tivemos bons momentos tanto dentro como fora do laborat´orio. Tamb´em sempre sou grato a Caroline Mouls que, mesmo distante, fez contribui¸c˜oes muito construtivas a esse trabalho.
mesmo em momentos de muita escurid˜ao e confus˜ao.
Lembro de agradecer tamb´em ao Laborat´orio Nacional de Nanotecnologia (LNNano) que disponibilizou a utiliza¸c˜ao de equipamentos necess´arios para a caracteriza¸c˜ao das amostras.
Lady Gaga, Dua Lipa e Sia por proporcionarem a trilha sonora.
Envio tamb´em votos de boa sorte a todos aqueles que, como eu, escolheram a carreira acadˆemica e est˜ao enfrentando as dificuldades que se apresentam atualmente neste cen´ario. Espero que vocˆes consigam realizar seus trabalhos da melhor maneira poss´ıvel.
Meu mais sincero muito obrigado `aqueles que lutam pela educa¸c˜ao num pa´ıs desgov-ernado!
Bem como devo agradecer a todos aqueles que, de alguma maneira, ajudaram na realiza¸c˜ao deste trabalho, mas que n˜ao foram mencionados anteriormente.
Igualmente sou grato aos ´org˜ao de fomento Coordena¸c˜ao de Aperfei¸coamento de Pes-soal de N´ıvel Superior (CAPES) pela bolsa de mestrado, processo 88887.145777/2017-00, Funda¸c˜ao de Apoio `a Pesquisa do Estado de S˜ao Paulo (FAPESP) e Conselho Nacional de Pesquisa (CNPq) pelo financiamento dos instrumentos utilizados durante este mestrado. Agrade¸co tamb´em ao Laborat´orio Nacional de Luz S´ıncrotron (LNLS) do Centro Na-cional de Pesquisas em Energia e Materiais (CNPEM) pela utiliza¸c˜ao das instala¸c˜oes e pelo programa de bolsas em conjunto com a CAPES e ao programa de p´os-gradua¸c˜ao do Instituto de F´ısica Gleb Wataghin que me permitiu obter este t´ıtulo. O presente trabalho foi realizado com apoio da Coordena¸c˜ao de Aperfei¸coamento de Pessoal de N´ıvel Superior - Brasil (CAPES) - C´odigo de Financiamento 001.
The main goal of this work is to obtain heterostructures of BaTiO3 (BTO) and Fe
with a magnetoelectric coupling at the interface. We can divide the results into two main lines: the optimization of the BTO layer and the growth of the bilayered structure used to characterize the coupling. In the first group, we were able to grow thin films by pulsed laser deposition (PLD), optimizing the growth parameters to obtain the best results. The use of different substrates is also a subject of analysis. We tried growing polycrystalline samples on silicon substrates. The results show that they could be crystallized after annealing in-situ up to 873 K. The morphology of the growth on SrTiO3 and Nb:SrTiO3
substrates was also studied. The BTO layer was grown very thick (∼ 200 nm) to ensure that the films would be less epitaxial. In the second group, we have grown the bottom electrodes of SrRuO3 by PLD while the titanium gold top electrode and the thin Fe
layer were grown by electron beam evaporation. The samples were characterized by piezoresponse force microscopy (PFM), showing that they were ferroelectric. The x-ray absorption spectroscopy (XAS) showed that the sample was metallic by the time it reached the microscope. While preliminary characterization by photoemission electron microscopy (PEEM) could indicate that there was a coupling between the ferromagnetic and ferroelectric layers.
Keywords: magnetoelectric coupling, multiferroics, photoemission electron microscopy, thin films.
O objetivo principal deste trabalho ´e obter heteroestruturas de BaTiO3 (BTO) e Fe
com um acoplamento magnetoel´etrico na interface. Podemos dividir os resultados em duas linhas principais: a optimiza¸c˜ao da camada de BTO e o crescimento da estrutura de mul-ticamadas usada para caracterizar o acoplamento. No primeiro grupo, pudemos crescer os filmes finos por deposi¸c˜ao por laser pulsado (PLD), optimizando os parˆametros de cresci-mento a fim de obter os melhores resultados. O uso de diferentes substratos tamb´em foi objeto de an´alise. N´os tentamos crescer amostras policristalinas em substratos de sil´ıcio. Os resultados mostram que elas puderam ser cristalizadas ap´os um tratamento t´ermico in-situ at´e 873 K. A morfologia do crescimento sobre substratos de SrTiO3 e Nb:SrTiO3
tamb´em foi estudada. A camada de BTO foi crescida bem espessa (∼ 200 nm) para garantir que os filmes fossem menos epitaxiais. No segundo grupo, crescemos eletrodos inferiores de SrRuO3 por PLD, enquanto eletrodos superiores de ouro titˆanio e uma
ca-mada fina de Fe foram crescidos por evapora¸c˜ao por feixe de el´etrons. As amostras foram caracterizadas por microscopia de for¸ca de piezorresposta (PFM), indicando que elas eram ferroel´etricas. A espectroscopia de absor¸c˜ao de raios X (XAS) apontou que as amostras eram met´alicas no momento em que entraram no microsc´opio. Enquanto a caracteriza¸c˜ao preliminar por microscopia eletrˆonica de fotoemiss˜ao (PEEM) pode mostrar ind´ıcios que h´a um acoplamento entre as camadas ferroel´etricas e ferromagn´eticas.
Palavras-chave: acoplamento magnetoel´etrico, multiferroicos, microscopia eletrˆonica de fotoemiss˜ao, filmes finos.
1 Introduction 12
1.1 Ferromagnetic domains . . . 15
1.2 Ferroelectric domains . . . 18
1.3 Magnetoelectric coupling . . . 23
1.4 Objectives of this work . . . 26
2 Experimental methods 28 2.1 Pulsed laser deposition . . . 28
2.2 Molecular beam epitaxy . . . 30
2.3 X-ray diffraction . . . 32
2.4 Atomic force microscopy . . . 33
2.4.1 Piezoresponse force microscopy . . . 35
2.4.2 Kelvin probe force microscopy . . . 37
2.5 Synchrotron-based techniques . . . 39
2.5.1 X-ray absorption spectroscopy . . . 39
2.5.2 X-ray magnetic circular dichroism . . . 41
2.5.3 Ferroelectric contrast . . . 41
2.6 Photoemission electron microscopy . . . 43
3.1.1 Group A: BaTiO3 on Si . . . 49
3.1.2 Group B: BaTiO3 on SrTiO3 and Nb:SrTiO3 . . . 51
3.1.3 Group C: BaTiO3 on SrRuO3/SrTiO3 . . . 53
3.2 In-situ transfer . . . 56
4 Conclusion and perspectives 66
Chapter 1
Introduction
The known ferroic orders are ferroelectricity, ferromagnetism, ferrotoroidicity and fer-roelasticity. When a material presents more than one of these orders, it is regarded as multiferroic. This definition was introduced by Schmid in 1994 [1]. A multiferroic material can involve any of the four mentioned ferroic orders. In this work, we are going to concen-trate only in materials or devices that are ferroelectric and ferromagnetic, thus, along this manuscript, the mention to a multiferroic will mean a ferroelectric-magnetic multiferroic. In the case that the multiferroic material presents a ferroelectric and a magnetic order, it can present a magnetoelectric effect, where the modification of one of the ferroic orders can modify the second. It is possible to trace the search of this magnetoelectric effect to the 19th century.
In 1894, Pierre Curie hypothesized about such an effect. It was only in 1926 that Debye coined the definition magnetoelectric effect. The pursuit of a material that presented this coupling remained fruitless for many years to come. In 1958, Landau and Lifshitz stated that this phenomenon was unworthy of discussion for it was not observed in any material [2]. Shortly after came the theoretical prediction [3] and experimental observation [4] of the magnetoelectric effect in Cr2O3.
than one ferroic order in a single phase. Therefore, they are difficult to be found in nature, since the ferroelectric and ferromagnetic orders are to a certain extent mutually exclusive. This is due to the opposing conditions for the d-shell of ferroelectrics and ferromagnets. Partially filled electron shells are required for the appearance of magnetic ordering, such as ferromagnetism. In the other hand, ferroelectricity can arise from several mechanisms. In some 3d transition metal based perovskites, including BaTiO3, it is know that the 3d0
occupancy causes an off-center shift of the transition metal, which forms strong ionic bonds with the oxygens. Empirically it was observed that any d-shell occupancy inhibits this process that creates a local dipole [5]. Therefore, ferrolectric materials require an empty d-shell [6], which usually precludes the presence of magnetism on the same material. The solution is to use the extrinsic multiferroic materials that present a coupling between two ferroic orders in more than one phase, such as multilayered structures.
The choice of Fe and BaTiO3 (BTO) for the materials composing the heterostructure
was motivated by the vast knowledge about both of them separately, and good lattice match. Also, a theoretical study stating that their coupling was more complex than just strain coupling [8]. In this study, the authors calculated that the iron atoms at the interface would induce a magnetic moment on the oxygen and titanium of the topmost layer of BTO. The oxygen magnetic moment would have a very weak contribution due to the small overlap of the O 2p levels and the Fe 3d levels. But the strong hybridization between the Ti and Fe levels leads to an antiferromagnetic alignment between the top Fe layer and the induced moments at the interfacial Ti atoms.
A previous work studied ultrathin (0.1 to 2 nm) layers of Fe by e-beam evaporation on top of single-crystal (001) BTO substrates [9]. This study concluded that the growth mode is by islands and that the interface is sharp, without oxidation on the Fe layer. Furthermore, a change in the easy axis of magnetization with the increase in the Fe layer
Figure 1.1: Schematic showing the two different types of magnetoelectric multiferroic compounds: the intrinsic (left) and the extrinsic (right). The first type (right) contains in a single phase both ferroic orders, while the latter (left) has an interface where the two ferroic orders interact. By manipulating this interaction, it could be possible to have an electric-field control of the magnetization (or a magnetic-field control of the polarization) of the sample [7]. In the middle, we have some hysteresis loops showing the relation between the external conjugated fields (E and H) and the order parameters (P and M). For ferroelectricity, the electric field (E) interacts with the polarization (P). While for ferromagnetism, the magnetic field (H) influences the magnetization (M). In the third row, we have the interaction of the electric field with the magnetization and the magnetic field with the polarization.
thickness could be identified.
In another work, it has been grown 10 nm and 20 nm of Fe by molecular beam epitaxy, on single-crystal (001) BTO substrates at 300◦C [10]. After growth, the samples were cooled down through the ferroelectric transition (Tc = 120◦C) to form the a-c domains
(see section 1.2). This group was able to demonstrate by MOKE microscopy that there is a coupling between the ferroelectric and ferromagnetic domains of the heterostructure. The growth of Fe was also studied on top of BTO thin films. In [11], the authors have
(30 nm)/BaTiO3 (50 nm) on a (001) NdGaO3 substrate. The bilayer was grown by
pulsed laser deposition. They investigated the interface of Fe and BTO by EELS, and found that Fe oxidised at the interface. This was also observed by another group on a SrTiO3/BTO/Fe heterostructure [12]. In this case, the Fe layer (1-3 nm) was grown by
molecular beam epitaxy on a BTO layer grown by pulsed laser deposition on a (001) STO substrate. The Fe interfacial atoms oxidation was measured by XPS.
Our group has also already obtained results for this system. Caroline L. Mouls in her thesis reported that a magnetoelectric coupling could be observed by photoemission electron microscopy (PEEM) for a Fe layer (1.5 nm) on top of single-crystal (001) BTO [13].
In the next sections we will give more details about ferromagnetism, ferroelectricity and the interaction between these two ferroic phases. We conclude this chapter discussing the objectives of this work.
1.1
Ferromagnetic domains
One would expect that the most stable configuration for a ferromagnetic material is that it would be a single dipole. But that is not the case. An actual ferromagnetic material is divided into several domains. We can see in Figure 1.2 why that is the case. The division in several domains reduces the magnetostatic dipolar interaction which is energetically more stable than a single dipole.
These domains can be at first randomly oriented, thus there would be no net mag-netization on the bulk material. With the application of an external magnetic field, the domains orientation would be induced to rotate and be aligned with the magnetic field direction. This process is represented in Figure 1.3. A net magnetization would arise in the material. When the external field is removed, the domains can relax but in the case
Figure 1.2: Schematic representation showing how the process of dividing a magnet re-duces magnetostatic energy within the material [14].
of permanent magnets the net magnetization does not average to zero and the magnet retains some net magnetic moment. That is the definition of remanent magnetization, i. e., the remaining magnetization after switching off the external magnetic field.
Figure 1.3: Process of changing the magnetization of several domains by an external magnetic field H that is oriented in the upwards direction [14].
The curve of the magnetization M as function of the external field H is called a hysteresis loop and can be seen on Figure 1.4. When we first start magnetizing the material, the curve follows a different path until saturation, this stage of the process is irreversible. The saturation of the magnetization happens when the driving field is able to align all of the domains inside the magnet.
The transition between two different domains is known as a domain wall (DW), and it changes through rigid spin rotations. If the rotation plane is parallel to the wall, it is regarded as the N´eel type. On the other hand, if the rotation plane is perpendicular to the wall, the type of DW is Bloch. Both can be seen in Figure 1.5.
Figure 1.4: Hysteresis loop of the magnetization (M ) in function of the external field (H) [15]. One can notice the remanent magnetization (Mr) mentioned before and the
demagnetizing field (Hci) that drive the magnetization to zero.
Figure 1.5: (a) Domain wall (DW) between two different magnetic domains. (b) Bloch domain walls. (c) N´eel domain wall [16].
1.2
Ferroelectric domains
In the case of tetragonal distorted ferroelectric perovskites like BaTiO3, the
ferro-electric domains are intrinsically related with the ferroelastic ones. These last ones are determined by the crystal structure, and so are the ferroelectric domains. The ferroelas-tic domain walls are related to crystalline twinning and here we first address the process of crystal twinning, since they determine the ferroelectric domains. The intention is to understand the possible ferroelastic domains in BaTiO3 to further on investigate the
fer-roelectric structures.
Two macroscopic examples of crystal twinning can be seen in Figure 1.6. They can come in two different types: penetration twins and lamellar (or contact) twins. The first ones are characterized as interpenetrating individuals with an irregular surface that have a common twin axis. On the other hand, lamellar twins are defined as having a definite composition surface, known as twin plane, separating the individuals. The presence of twinning in a crystal can change the material properties, such as piezoelectricity, since different twin components contributions can cancel each other out.
A process that can be the cause of twinning in a crystal is mechanical twinning. It consists in the application of mechanical force favoring the development of certain twin components to detriment of others. The appearance of new and more suited twin components from single-crystal specimens is also a possibility. This has been known for metal elements for a long time, that when one part of the crystal would be ground against the other under shear stress [17].
In the scope of this work, to understand ferroic domain formation, we would like to focus on transformation twins. They are a type of crystal twins that can be described by two crystal structures SG and SF. These symmetry groups have two defining properties:
SF has lower symmetry than SG and it arises from small distortions of the SG lattice.
(a)
(b)
Figure 1.6: Real examples of twinned crystals. (a) As-grown contact twinned crystal of rutile TiO2 [17]. (b) Penetrating twinned crystal of pyrite [18].
TT R. An example (Figure 1.7) is a crystal with a tetragonal unit cell. This implies that
the transition from the higher symmetry group to the lower one shifts the ions about 10−2 - 10−3 nm around the equilibrium position in the c-axis. This change breaks some symmetries of the original group (SG) reaching a lower symmetry group (SF). In this
case, SF can have two orientations in relation to the dislocation sign of the central atom:
SA and SB. Without external interference, the formation of either is equally likely. These
two structures (SG and SF) are transformation twins. The twin components (SA and SB)
are called domains.
Figure 1.8 shows the specific case for BTO. It presents a phase transition from para-electric cubic phase (P m¯3m) to tetragonal ferroelectric phase (P 4mm) at 393 K. This phase is stable until 278 K, when the structure transforms into orthorhombic (Amm2). The final transition happens at 183 K into a rhombohedral symmetry (R3m) [19]. The Ti atoms dislocation to one direction, for example, “upward”, is followed by the dislocation
Figure 1.7: Simple example of ferroelectric domain formation on a crystal lattice [17].
Figure 1.8: Schematic of the perovskite crystal structure of BaTiO3. Paraelectric cubic
phase (A). Ferroelectric tetragonal phase (B and C). The dislocations are exaggerated for clarity [19].
of the two different types of oxygen sites in the “downward” direction [17]. These dis-locations create an electric dipole. Thus, we can see the polarization in BTO is directly correlated to the crystalline structure.
Once again, we can have a region between two domains called the domain wall (DW). In the case of electric polarization, since it is not quantized, it can continuously change in amplitude inside the DW. This implies that a different type of DW than the usual two already discussed for magnetism is possible: the Ising DW. This DW is shown in Figure 1.9 along with the Bloch and N´eel ones. This type of DW is energetically more favorable in relation to the chiral ones, for two reasons: the piezoelectric coupling between
Figure 1.9: (a) Ising wall, (b) Bloch wall, (c) N´eel wall and (d) mixed Ising-N´eel wall [20].
the polarization and spontaneous strain makes the rotation energetically unfavorable. Moreover, there is also an electrostatic cost, since any change in polarization away from the wall would induce an accumulation of charge in the walls. For 180◦domains the Ising DW was seen as the most common, until recently when it was shown that chirality can be induced at the DW, due to the multiferroic coupling. Such is the case for perovskite ferroelectrics, since they are ferroelectric and also ferroelastic [20].
In the following we discuss the case of a tetragonal BaTiO3 single crystal. We adopted
the usual nomenclature for quasi cubic perovskite ferroelectric domains, in which the axis ’a’, ’b’ and ’c are respectively parallel to the [100], [010] and [001] crystalline directions of the quasi-cubic structure. In each domain the ferroelectric polarization must be along the tetragonal axis, which in the case of BaTiO3 is 0.0044 nm longer than the other two
Figure 1.10: Possible orientations of the ferroelectric phase (inner circle) and polydomain structures (outer circle) [21].
(a)
(b) (c)
Figure 1.11: Different types of domains that can be present in BTO. (a) 90◦walls between a1− a2 domains. (b) 180◦walls. (c) 90◦and 180◦walls forming the herring bone structure
[22].
two). Thus, the tetragonal axis can point along any of the ’a’, ’b’ and ’c axis. This can be seen in Figure 1.10. Considering that the ferroelectric polarization can be only parallel or antiparallel to the tetragonal axis, we end up with 6 possible ferroelectric domains, namely, the polarization can assume only two directions, always parallel to one of the three axis of the pseudo-cubic strucutre. In Figure 1.11, we can see the different types of
is possible to have c − a domains and a1− a2 domains with 90◦ and 180◦ walls. Here we
defined as c the out of the plane axis.
1.3
Magnetoelectric coupling
A simple approach to understand the magnetoelectric coupling is to expand the free energy in relation to the electric and magnetic field. Details of such derivation can be found in reference [23] and eq. 2.17 of [24]. In this way, the free energy (F ) reads as:
F ( ~E, ~H) = F0− PiSEi− MiSHi− 1 2χ e ijEiEj− 1 2χ m ijHiHj− αmeij EiHj − 1 2βijkEiHjHk− . . . (1.1) where F0 is a constant term, independent of the ~E and ~H fields, PiS and MiS are the
spontaneous polarization and magnetization respectively, χe and χm the corresponding electric and magnetic susceptibilities and αmeand β are higher order tensors that mediate the magnetoelectric effect. The ij indices indicate that the .Noting that χe
ij = 0ij and
χm
ij = µ0µij (where 0 and µ0 are the vacuum electric and magnetic permeabilities and
ij and µij are the medium electric and magnetic permeabilities). As the permeabilities
can be anisotropic, they are represented by three dimensional tensors, where i and j are the three possible ortogonal directions. We can obtain the total polarization and magnetization by taking the derivatives of the free energy with respect to the electric field and magnetic field:
Pi( ~E, ~H) = −
∂F ∂Ei
Mi( ~E, ~H) = −
∂F ∂Hi
= MiS+ µ0µijHj+ αijEj (1.3)
We can notice that the expressions for the total magnetization and polarization are composed of three parts: the spontaneous one (PiS and MiS), the induced by a conju-gated field (0ijEj and µ0µijHj) and the magnetoelectric one (αijHj and αijEj). In this
analysis, higher order terms are being neglected. The coefficient αij is the linear
mag-netoelectric tensor, a second rank axial tensor. If we can draw conclusions about it, we are getting insight into the magnetoelectric coupling. Symmetry considerations can give information about the availability of magnetoelectric materials. But they do not account for the strength of the coupling. In that regard, a very important result that was shown is that this magnetoelectric tensor is upper bounded [25]:
α2ij < χeijχmij (1.4)
Therefore, only materials that present high susceptibilities are allowed to have a large magnetoelectric effect. Since ferromagnetic (ferroelectric) materials have high magnetic (electric) susceptibilities, materials that show both ferroic orders in the same phase are the most searched to have a large magnetoelectric coupling. However, given the previously mentioned difficulty in finding single phase materials that present ferroelectricity and ferromagnetism, when one of the susceptibilities has a high value, the other has a low one, making the magnetoelectric coefficient limited. The reason for this search is related to the possible applications of such material. It can be used to develop a magnetoelectric memory [26] or to deliver drugs with magnetoelectric nanoparticles [27].
To overcome the limitation on the magnetoelectric coupling coefficient of single phase materials, one approach is to use two phase systems (artificial multiferroic materials) that would present at the interface an indirect coupling. This kind of materials could present magnetoelectric coefficients orders of magnitude higher than the intrinsic ones
Figure 1.12: Images taken by photoemission electro microscopy (PEEM) with two different contrasts. To the left, the ferroelectric contrast of the domains from the BaTiO3
single-crystal. To the right, we have the ferromagnetic domains of the Fe thin film given by X-ray magnetic circular dichroism (XMCD). It is possible to see the coupling of the two layers since the domains have the same stripe pattern, expected for in-plane (a1 − a2)
ferroelectric domains. Adapted from [13].
[28]. Their coupling could be strain-mediated. An electric external field would change the size of the ferroelectric lattice parameter due to the converse piezoelectric effect. This change would be transferred by contact to the magnetic material and change its magnetic properties, there is, the anisotropy by magnetostriction [29]. This is know as the converse magnetoelectric (ME) coupling and it is defined as the product of the magnetostrictive effect of the ferromagnetic material with the piezoelectric effect of the ferroelectric one.
The other form of indirect coupling is the charge-mediated. This is caused by the interaction of the ferroelectric polarization direction with the electronic structure of the ferromagnetic layer. Several mechanisms are responsible for this electronically-driven change in magnetic properties. The first one is the spin-dependent screening in the fer-romagnet of the interface-bound charges of the ferroelectric [29]. Since the spin up and
down electronic density is different at the Fermi level of ferromagnets, the screening is spin-dependent. Therefore, this leads to changes in the surface magnetization and magne-tocrystalline anisotropy. The second one is a change in electronic bonding at the interface between the ferromagnet and the ferroelectric [29]. The displacement of the ferroelec-tric atoms due to a polarization reversal changes the overlap between the orbitals of the interfacial atoms. This leads to a difference in charge transfer that alters the interface magnetization, anisotropy and spin polarization.
This brings us to the system at hand: Fe thin films on top of BaTiO3 (BTO). As we
have already mentioned, the coupling at the interface may be more complex than just a strain interaction. Figure 1.12 shows an example of the magnetoelectric coupling for Fe thin film on top of BTO single-crystal. The methods used to obtain the contrast in this image will be discussed in chapter 2. Here we point out that the left part of the figure shows the stripe-like ferroelectric domains of a BTO monocristal. These stripes are typical of a1 − a2 domains, in other words, ferroelectric polarization lying at two
orthogonal axis in the plane of the monocrystal. At the right of the same figure one can observe the magnetic domains of a thin Fe layer deposited on the top of the monocrystal. We would expect that the thin Fe layer would have randomly oriented domains. But one can notice that the magnetic domains follow the ferrolectric stripes, which are a direct way of determining the existence of a magnetoelectric coupling between the ferroic orders.
1.4
Objectives of this work
As mentioned before, several works have been published on thin Fe films on top of BaTiO3 (BTO) bulk single crystals [8–10], because in this case the domains are large,
usually several micrometers wide by hundreds of micrometers long. This allows one to use optical techniques in their investigation (e.g. [10]). But for applications the use of BTO thin films is mandatory instead of single-crystals. Normally BaTiO3 thin films have
and ferromagnetic domains if the ferroelectric domains have in-plane orientation. It is expected that out-of-plane ferroelectric domains would also be coupled with the ferro-magnetic domains, but here we concentrate on in-plane ferroelectric and ferroferro-magnetic domains. We have chosen this approach so that one could compare the results obtained for BaTiO3 single crystals, which present macroscopic domain sizes, with those eventually
obtained in BaTiO3 thin films, which normally present domains in the nanometer range.
For that reason we need to grow BTO films with in-plane ferroelectric domais. It happens that in our setup the deposition system is isolated from the microscope. Since we want to study the system Fe on top of BTO, we have to guarantee that the Fe stays metallic, that is, without oxidation. We could use a capping layer, but that would limit the PEEM technique. Therefore, we need to develop a vacuum suitcase-like sample transfer system between the growth and PEEM chambers under ultra-high vacuum (UHV) conditions.
Thus, this work has two main objectives:
• Find the optimal growth parameters for Fe and BaTiO3 bilayer thin films with
in-plane ferroelectric domains
• Develop an instrumentation enabling in-vacuum transfer from the growth chambers of the pulsed laser deposition (PLD) and molecular beam epitaxy (MBE) systems to the photoemission electron microscope (PEEM)
In the following sections, we are going to further describe the work. In chapter 2, there is a complete description of the growth and characterization techniques employed. The samples characterization is presented in chapter 3. And finally, we conclude the work in chapter 4.
Chapter 2
Experimental methods
In this chapter, we are going to describe the techniques and methods used to grow and characterize the samples. The samples were grown using pulsed laser deposition (PLD) and molecular beam epitaxy (MBE). The characterization was made with X-ray diffraction (XRD), atomic force microscopy (AFM), piezoresponse force microscopy (PFM), Kelvin probe force microscopy (KPFM) and photoemission electron microscopy (PEEM). We are also going to describe the process of generating synchrotron light and detail the techniques that used it: X-ray absorption spectroscopy (XAS), X-ray magnetic circular dichroism (XMCD) and X-ray linear dichroism (XLD).
2.1
Pulsed laser deposition
As in this work, we planned to grow BaTiO3 (BTO) thin films, pulsed laser deposition
(PLD) was the growth technique of choice. PLD has a great advantage as it keeps the target stoichiometry in general. This is particularly important for complex oxides, like BTO. Indeed, part of the advantage of PLD comes from the fact that it allows depositions under quite high oxygen pressures in the chamber, which greatly limits the presence of oxygen vacancies in the films. Furthermore, it can achieve higher growth rates under
frequency magnetron sputtering.
Figure 2.1: Schematic of the system for pulsed laser deposition. The inset shows the plume. [31]
The basic setup of this technique can be seen in Figure 2.1. The deposition consists in a target ablation by an ultraviolet laser. This process creates a plume of ejected material (shown in the inset of Figure 2.1). The plume is highly directional and has a distribution of mass and velocity of the ejected material. This means that by controlling the distance between the substrate and the targets, one can select the deposited material energy and the size, tuning it to the desired application. The parameters that can be controlled during the deposition are the main chamber atmosphere (which gases and pressure), the substrate temperature, the laser fluency and the distance between the substrate and the target. The fluency is defined by the energy distributed on an area by the laser. By controlling the laser energy per pulse or the spot size, it is possible to change this parameter. A system
of focusing lenses can also change the spot size.
Several kinds of laser can be used for ablation, the basic requirement is being able to deliver nanosecond pulses of tenths of joules in energy. These are normally the parameters needed to obtain a controllable laser ablation. The system we used is equipped with a Coherent Compex Pro 102 F excimer laser, operating with Kr based mixture as exciting gas. This laser delivers pulses of up to 400 mJ, in 20 ns, with the maximum of 20 Hz repetition. The deposition chamber is from the company SPECS GmbH.
2.2
Molecular beam epitaxy
We will use this technique to grow approximately 1 nm of iron on top of the BaTiO3
thin films. We have chosen this technique because it can grow high quality films without oxidation. This would be very difficult to do inside the PLD chamber since the base pressure has always oxygen contamination.
This growth technique is mainly used for semiconductor thin films. It is very versatile and clean to grow simple or complex films under ultra-high vacuum (UHV) or reactive atmospheres. The basic setup for this technique is shown in Figure 2.2. It works by heating the material to be deposited inside effusion cells. The heating is made by an electron beam. As the material heats up, it reaches its vapor pressure and begins to evaporate. Since the cell has an opening in the substrate direction, a collimated jet of atoms is ejected in the direction of the substrate, hence the name molecular beam epitaxy. As this beam of atoms reaches the substrate surface, which can also be heated, it begins to rearrange to grow the film. The atoms flux can be controlled by a shutter in front of the cells. It is possible to grow several materials at once. Some methods can be used to monitor the film growth rate, such as quartz crystal microbalance (QCM) or reflection high energy electron diffraction (RHEED). We have both methods available in our chamber.
Figure 2.2: Simple schematic showing how molecular beam epitaxy works. [32]
There are three main zones depicted in the image: the molecular beam generation zone, the vapor elements mixing zone and the substrate crystallization zone. In the first one is where there is the formation of the molecular beam as described before. Then when there are more than one material being evaporated or if there is an reactive atmosphere, the mixing of the vapor elements happens in the second zone. In the third zone, it is where the deposition occurs with the arrival of the mixed elements on the surface and the reorganization of the deposited material with help of the energy provided by the heating block. This is the process of crystallization.
For growing the Fe thin layer, we used the QCM to find the desired parameters to grow approximately 1 nm of Fe. The vacuum inside the chamber was 1.51 10−9 mbar during growth.
2.3
X-ray diffraction
We used X-ray diffraction to evaluate the crystalline quality of the grown films and to check if the film has grown as expected. We can also get information about the thickness and in-plane strain of the film. It is also possible to probe the epitaxy between the film and the substrate.
X-ray diffraction has been used as a structural technique since the beginning of the 20th century. The mechanism of Bragg’s law that explains the phenomenon is very well known. One can see a simple explanation in Figure 2.3. Each row represents a plane of atoms. Diffraction patterns are the interference from rays scattered on the material surface with the rays scattered by planes of atoms inside the material. The difference in the path between these two rays is calculated and if it equals a multiple of the rays wavelength, then the interference will be constructive and we will have a peak for this incidence angle. If we change the incidence angle, we can map the sample entire interference pattern. This is called a θ - 2θ measurement.
Figure 2.3: Explanation of Bragg’s law. [33]
This type of measurement is called X-ray reflectivity (XRR). Analyzing the interference fringes it is possible to obtain information about the thickness of samples, interfaces roughness, and electron density for each layer of the sample. In this work we limited the analysis to obtain the thickness of the layers.
2.4
Atomic force microscopy
Atomic force microscopy (AFM) is based on the force between a sharp tip and the surface to be probed. When the tip is at a few nanometers of the surface it feels an attractive van de Waals force, while if it approaches too much the force starts to be repulsive. One operation mode of AFM microscopes is to feed the force signal to a control electronic loop in order to keep the distance between tip and surface constant. Scanning the tip across the surface and measuring the tip position needed to keep it at constant force allows one to obtain an image that represents the topography of the sample. There are other possible sources of forces involved in the tip – surface interaction that can also be used as feedback, given other kind of information about the sample surface. Some review works about AFM based techniques can be found in references [34, 35]. In the following we give more details about AFM directly related with our results.
We are interested in using atomic force microscopy (AFM) to evaluate the topography of the samples. Using this technique, it is possible to see what is the growth mode of the thin film: by islands, layer-by-layer or if it begins layer-by-layer and turns into islands afterwards. We can obtains values for the surface roughness evaluating the flatness of the sample surface.
The basic setup for an AFM measurement is shown in Figure 2.4. The sample is placed on a piezoelectric base that can move on the three directions. A laser shines on the back of a very thin cantilever with a tip. Therefore, the cantilever movement can be
Figure 2.4: Schematic of an atomic force microscope. [36]
measured by analyzing the signal on each quadrant. This can be better seen in Figure 2.5. There are two modes of operation for this microscopy. The first is contact mode, which keeps the tip in contact with the sample surface. The other is dynamic (tapping) mode, which the sample oscillates in intermittent contact with the surface. Both modes are controlled by a feedback loop, which is the basis for a PID controller.
The microscope used was a Nanosurf FlexAFM. The tips used were Multi75E-G for the electrical measurements and Tap300Al-G for the AFM ones.
Figure 2.5: Cantilever movements and optical deflections. [37]
2.4.1
Piezoresponse force microscopy
We have used this technique, because it enables direct measurement of the surface ferroelectric domains of the sample. This is important to be used in comparison with photoemission electron microscopy images of the ferroelectric domains. It is also possible to use the same setup to write ferroelectric domains applying a strong electric field on the tip. This will be better described in this section. By writing domains, we can make sure that the sample is ferroelectric and that the top and bottom electrodes are working.
Piezoresponse force microscopy (PFM) follows the same principle as the AFM, the difference is the use of a conductive tip. The measurement is conducted in the contact mode. We apply an AC bias to the tip and measure how the surface responds to this electric field. The response amplitude is proportional to the magnitude of the local elec-tromechanical coupling, while the phase provides the local polarization orientation. The phase relation with the polarization is explained in Figure 2.6. When the tip electric
Figure 2.6: Representation of how the piezoresponse of the material is related to the ferroelectric domains orientation. [37]
field is in the same sense of the surface polarization, the response is in phase with the excitation. The reverse happens when they are anti-parallel, causing a 180◦ phase shift in the response.
One can also use the electrical tip in contact mode to polarize the sample. This process is done by applying a DC-voltage between the tip and the bottom electrode. This voltage can be varied from -20 V to 20 V. It can be used to see if the sample maintains the polarization after the electrical excitation is retreated. This would indicate that the sample is ferroelectric. Another application would be making patterned domains on the sample.
With this technique, it is possible to measure the surface local work function of the sample and this is directly correlated with ferroelectric domains. Out-of-plane domains have a charge accumulation due to the polarization dipole that changes the local work function. Therefore it is also possible to have an insight into the domain configuration of the sample using this technique.
Figure 2.7: Schematic of a measurement with a kelvin probe force microscope. 1) An AC bias with frequency ω is applied between the tip and the surface. 2) An electromechanical oscillation with frequency ω is detected by the photodiode. 3) This signal is given to the feedback loop that tries to nullify this contribution by applying a tip voltage. [38]
This technique also uses a conductive tip, but it is measured on the tapping mode. An AC bias is applied to the oscillating tip which creates an oscillating capacitance between the tip and the surface. The interacting potential between the tip and the sample is given by:
∆V = ∆φ − Vdc+ Vacsin(ωt), (2.1)
ampli-tude and ω is the applied AC voltage signal frequency. The energy of a parallel plate capacitor is given by:
U = C∆V
2
2 , (2.2)
where C is the local capacitance between the sample and the tip. The force between the tip and the sample will be given by:
F = −∂U ∂z = − 1 2 ∂C ∂z∆V 2 = F dc+ Fω+ F2ω (2.3) ,
where the three components can be analyzed separately. The DC component is:
Fdc = − 1 2 ∂C ∂z " (∆φ − Vdc)2+ Vac2 2 # (2.4)
And the two components at frequencies ω and 2ω:
Fω = − ∂C ∂z [(∆φ − Vdc)Vacsin(ωt)] (2.5) F2ω = 1 4 ∂C ∂z[V 2 accos(2ωt)] (2.6)
If the component Fω is zero, then we have ∆φ = Vdcand in this manner, we obtain ∆φ
locally by plotting the DC tip voltage. By calibrating the tip work function, it is possible to determine the surface work function.
The 2ω component (second harmonic) is directly proportional to the sample local dielectric properties (∂C/∂z). This can bring insight into the sample surface defects and heterogeneities.
Synchrotron radiation is caused by the emission of relativistic electrons on circular paths. Every time the electrons are accelerated perpendicularly to their path, they emit a broadband tangential radiation with high brilliance. This radiation can be collected by experimental stations known as beamlines that use it for different applications. Histori-cally, the radiation was produced with the help of bending magnets. Nowadays, insertion devices are implemented to produce brighter energy beams. These are undulators and wigglers. Both are based on producing a sinusoidal back and forth motion of the elec-trons to produce radiation. Wigglers make this motion very wide so the individual rays of radiation produced by each wiggle cannot interfere with one another. Undulators, on the other hand, are built to use this interference to produce higher-order harmonics.
2.5.1
X-ray absorption spectroscopy
We will use this technique to obtain direct insight about the state of oxidation of the Fe thin layer. It will be also used in conjuction with the photoemission electron microscopy (PEEM) to obtain the different contrasts on the images, both ferroelectric and ferromagnetic.
X-ray absorption spectroscopy (XAS) obtains information on the atoms unoccupied states. The energy of incident X-rays must be variable for this technique to be employed. Therefore, it became more popular with the creation of synchrotron light facilities. There are two ways to obtain the spectra: the direct way and the total electron yield (TEY). In the first one, we measure the beam intensity before and after the sample. The difference in intensity is related to the absorbed total. This method is the most used with hard X-rays since they are more energetic and can penetrate greater distances. Softer X-rays cannot go through the samples. In this case, TEY is the method to be used. The sample is grounded, allowing the flow of a current to compensate the electrons lost. This current
is measured and normalized by the current of a gold screen which receives the beam before the sample. The sum of all the electrons ejected (the current) is proportional to the absorbed radiation, yielding the same result. The absorption spectra are sensitive to the chemical element, including its oxidation state.
The quantum description of this technique is based on the Fermi’s Golden Rule, which uses the time dependent perturbation theory to calculate the transition rate wif from an
initial state |ψii with energy Ei to a final state |ψfi with energy Ef [39]:
wif = 2π ¯ h hψf|Hint|ψii + X k hψf|Hint|ψki hψk|Hint|ψii Ei− Ek− hν 2 δ(Ef − Ei− hν) (2.7)
The delta term makes sure that the energy is conserved as given by the photoemission equation:
Ek = hν − Eb− φ, (2.8)
where Ek(or Ef) is the kinetic energy of the photoelectrons, hν is the photon energy, φ
is the work funtion of the material, Eb is the binding energy of the electron (Ei = Eb+ φ).
The interaction Hamiltonian is given by:
Hint= X j − e mcA(rj) · pj+ e2 2mc2A 2(r j) ! , (2.9)
where A is the potential vector operator from the electromagnetic wave, p is the electron momentum operator, rj is the electron position operator and the sum is taken
for every electron of the system.
These equations, along with approximations of the incident photons energy and of the states of the other electrons in relation to the emitted one, allow the calculation of the material absorption coefficient. This result can be directly used to analyse the XAS data.
We will use this technique to obtain the ferromagnetic contrast on the PEEM images. It is heavily based on XAS. The difference is that two absorption spectra are taken with different circular polarization. This can be seen in Figure 2.8. If we subtract one spectrum from the other, we obtain the signal from the X-ray magnetic circular dichroism (XMCD). This signal is proportional to the sample magnetism. This happens because the d levels occupation of the transition metal atoms have a difference between the spin up and spin down electrons. Then the absorption changes from one edge to the other, as it is represented in Figure 2.9.
Figure 2.8: Absorption spectra for the L2 and L3 iron edges with the two circular polar-ization. The blue signal is the difference between the spectra.
2.5.3
Ferroelectric contrast
Similarly to circular dichroism, this technique uses the linear polarization of light to obtain different spectra that give information about the sample. In the scope of this work, we will be using this technique to probe the ferroelectric domains of BTO. As can be seen in Figure 2.10, the d orbitals of the Ti split into two orbitals eg (dx2−y2 and dz2) and t2g (dxy, dyz and dxz) because of the hybridization with the oxygen atoms in BTO. During
Figure 2.9: Schematic XMCD effect in a single electron picture with large exchange-splitting in the d-bands. Transitions are occurring from the 2p spin-orbit split levels to the 3d by excitation from circularly polarized light. The probabilities are calculated by using the Clebsch-Gordon coefficients by taking into account the angular momentum of the photons. [40]
the ferroelectric phase, the Ti atom dislocation is in one of the three main directions x, y or z coming closer to one oxygen atom. This causes the orbitals in the plane perpendicular to this direction to have a lower energy. Thus, for direction z, this orbital is dxy. As for
direction x (y), this orbital would be dyz (dxz). The split between eg and t2g can be seen
on the absorption spectrum, but the split between these orbitals is not visible.
If the photons polarization is aligned with the direction of one of these orbitals, it causes the absorption to differ. A difference in absorption gives a difference in electrons being ejected, which affects the PEEM image contrast. The contrast is independent of the domains direction, that is, if they have polarization to the positive or negative z direction, for example. Therefore, they are only sensitive to the polarization axis. We can increase the contrast by two methods: either subtracting images with the same energy and different linear polarization or images with different energies and the same linear
Figure 2.10: Splitting of the Ti d orbitals for a BTO domain in the z direction. [13]
polarization. These energies have to be chosen to break the symmetry inside the t2g
orbital.
2.6
Photoemission electron microscopy
Using this technique, it is possible to produce images of the ferroelectric and ferro-magnetic domains of the sample. It will be used to evaluate if there is an interaction between the ferroelctric domains of the BaTiO3 layer with the ferromagnetic domains of
the Fe layer.
Photoemission electron microscopy (PEEM) has a basic setup represented in Figure 2.11. We shine the sample with ultraviolet or X-ray light. The electrons from the sample are excited and are ejected from the atoms. These electrons have very characteristic
kinetic energy given by the photoemission equation 2.8. They generate a cascade of secondary electrons either by inelastic collisions with the lattice atoms or by relaxation of the holes they left behind. The total current of electrons leaving the sample is collected by the objective lens and accelerated towards the prism. The role of the prism is to focus the electrons as well as to deflect their path towards the projector lenses. The projector lenses are very similar to the column of a conventional electron microscope. At the end of the column, the electrons are multiplied by a multichannel plate (MCP) and projected on a phosphor screen. A CCD camera captures the image on this screen.
Figure 2.11: Schematic of the photoemission electron microscope from SPECS. [41]
chemical or magnetic contrast depending on the interest of the application. For example, a magnetic image can be obtained by taking two images at different photon circular polarization (right and left) and subtracting one from the other. This uses the principles of X-ray magnetic circular dichroism (XMCD) to create an image contrast. Furthermore, it is also possible to use beams of different energies to chemically select the elements from a sample. Sample regions will emit electrons at different rates if illuminated by light with energy close to the absorption edges of the elements that are present in these regions. Absorption spectra details can also be analyzed in the same manner.
2.6.1
XMCD-PEEM
Figure 2.12: Process of obtaining the PEEM image with magnetic contrast. First, we measure what we call a quartet which consists of four images with polarization in the pattern plus, minus, minus, plus. Then we take the average between the two plus (blue line) and the two minus (orange line) images. Finally, we subtract one average from the other to obtain the image with magnetic contrast.
In Figure 2.12, we can see the process to obtain a PEEM image with XMCD contrast. We take four images with different polarization. This is what we call a quartet. The
The pattern is plus, minus, minus, plus. After that, we take the average from each group of polarization and subtract one from the other, much like we do with spectra. This results in an image with magnetic contrast.
Chapter 3
Results
As mentioned in section 1.4, this Master work has two main objectives, namely: i) obtain thin BaTiO3 films by PLD with ferroelectric domains not only out of the plane as usual, but also in the plane (“a1− a2” type of domains); ii) to develop a vacuum suit case
and a sample-holder to integrate the photoemission microscope (PEEM) with the thin-film growth system available at the LNLS, avoiding to expose the samples to atmosphere. We want to avoid the exposure to ensure that the Fe thin layer stays metallic. Thus, this chapter is divided into two main sections. Initially we will discuss the results obtained for the thin ferroelectric BaTiO3 film. In the final section we discuss the results obtained
for the instrumentation developed along this work.
3.1
Thin film growth
As mentioned in the introduction, we are looking for BaTiO3 films with in plane
ferroelectric domains. We have taken two approaches to achieve this. Firstly we tried to grow on a amorphous substrate to impose no preferential crystalline orientation on the sample growth (Group A). The substrate choice was SiO2/Si. The samples were grown
grow thicker samples to allow the BTO layer to relax enough to develop a polycrystalline structure (Group B). The thickness was estimated to be around 500 nm. Lastly, we have grown moderately thick samples (∼ 100 nm) with a SrRuO3 and a top electrode of
titanium gold (Group C). Thus, the samples are divided in these three groups.
We have grown 20 samples in total. This section is dedicated to present the results obtained for a selected group of grown samples. An important part of this work was to optimize the PLD growth parameters for BaTiO3. We started from values found in the
literature (e.g., [11, 12]). In the following sections we will describe the growth parameters for each sample, but at the end several of those are the same for all samples and we list them below with a brief description of their influence on the thin film growth:
• Growth oxygen pressure: one normally wants this pressure to be as high as possible to be sure that the sample does not have oxygen deficiency. On the other hand, the plume characteristics strongly depend on this pressure, and as it increases the deposition rate decreases. In our case we kept the oxygen pressure during deposition at 0,1 mbar.
• Annealing oxygen pressure: this pressure is used to preserve the stoichiometry of the film during the crystallization process in the annealing. An annealing under vacuum would result in oxygen vacancies. We kept this pressure at 100 mbar.
• Target-sample distance: the size of the plume is different depending on the pressure of the chamber. By changing this distance, we can select different regions of the plume, changing the growth rate and the size and energy distribution of the ejected material. Having selected an oxygen pressure, we must tune this distance to avoid splashing (phenomenon that occurs when big chunks of the ejected material are directly deposited in the substrate) and maximize the growth rate. We have used 50 mm as this distance.
This can have an influence in the roughness of the films and also change the process of crystallization during growth. We have kept this parameter at 3 Hz.
Those above are the parameters that were kept constant for all samples. Besides them, we should mention the following that will be quantified for each set of samples:
• Growth temperature: to guarantee during growth that the material reaching the substrate has enough energy to diffuse and begin the crystallization process, we provide the substrate with thermal energy. This has to be high enough that the crystallization can occur, but not so high as to re-evaporate the material.
• Annealing temperature: following the same logic as the previous parameter, we provide thermal energy so that the crystallization process post growth can happen during the annealing.
• Number of laser pulses: this is directly related to the amount of material ejected, therefore it will be proportional to the thickness of the film. The film will be thicker if we increase the number of laser pulses.
• Laser fluency: this is another parameter that is directly related to the plume char-acteristics. It is the rate of energy per area of the laser spot hitting the target. By changing it, we can also control the size of the plume and the distribution of ejected material through it.
3.1.1
Group A: BaTiO
3on Si
Here we present the result for one film of BaTiO3 on Si substrates. A few of these
were fabricated, but this is the sample that presented the best results. We recall that the intention here was to use the intrinsic amorphous SiO2 top layer to force the film to be
Figure 3.1: Diffraction pattern showing the crystallization of a sample grown on a Si substrate. The marked regions highlight the differences in the diffractograms as temper-ature increases. This are due to the crystallization of the SiO2 layer of the substrate.
The most important change is the appearance of the (101), (110) diffraction peak of the polycrystalline barium titanate.
The growth parameters of this sample are: substrate SiO2/Si (001), laser fluency of
1,3 J/cm2, 10000 laser shots, growth temperature of 201◦C, annealing temperature of 203◦C and annealing time of half hour.
In Figure 3.1, we can see the process of crystallization of the sample grown on a silicon substrate. The temperature was slowly increased as we measured the different diffraction patterns. This was made at the XPD (X-ray Powder Diffraction) beamline at UVX-LNLS, with 10 keV energy X-rays, using θ − 2θ geometry. The in-situ annealing of the samples were done by means of a resistive furnace attached to the diffractometer. The main feature is the formation of the BTO (101), (110) reflection when at approximately 600◦C. This is the most intense peak of a polycrystalline BTO thin film.
Figure 3.2 shows the diffraction patterns before and after the in-situ annealing. It reveals that the sample initially did not have a polycrystalline structure and after the
Figure 3.2: Diffractogram of the sample grown on Si substrate before and after the in-situ annealing. The peaks marked with a * are due to the SiO2 substrate. The measurements
were made in a diffraction beamline at 10 keV energy.
These results show that we are able to grow polycrystalline BaTiO3 on SiO2/Si
sub-strates. We could also observe the crystallization of the amorphous SiO2 layer during the
annealing ex-situ at the diffraction beamline (regions marked by the boxes in Figure 3.1). We did not pursued this option during this work, because given that SiO2 is insulating,
we would need to grow an electrode layer under the BTO to characterize it as a ferroelec-tric. We preferred to try the more standard SrRuO3 on SrTiO3 solution as electrode (see
section 3.1.3). But the possibility of growing BTO directly on SiO2/Si is very interesting
on the point of view of applications as will be discussed in the conclusions chapter.
3.1.2
Group B: BaTiO
3on SrTiO
3and Nb:SrTiO
3We present the results for two samples in this group. Their growth parameters are: laser fluency of 1,3 J/cm2, 50000 laser shots, growth temperature of 770◦C, annealing temperature of 655◦C and annealing time of 2 hours. Samples B1 was grown on top of SrTiO3 (001) and B2 was grown on Nb:SrTiO3 (001).
Figure 3.3: PFM (red) and KPFM (blue) images of the sample grown on STO (B1). a) and d) show the topography signal, the first is taken unsing contact mode for PFM and the second using tapping mode for KPFM, b) is the piezoresponse signal phase and c) is the lateral piezoresponse signal phase. e) is the tip voltage and f) is the second harmonic. The grain size is approximately 50 nm.
The topographic measurements made can be seen in Figures 3.3 and 3.4 for the sam-ples B1 and B2, respectively. One example of each substrate is shown to compare the similarities and differences.
For sample B1, the surface granularity is very small (∼ 50 nm). The PFM measure-ments, Fig. 3.3 - b) and c), show that the piezoresponse is localized on the small grains, indicated by the light/dark contrast in the images. There are differences between the in-plane (Fig. 3.3 - b) and out-of-plane (Fig. 3.3 - c) components for a single grain. We interpret this as a sign of ferroelectricity, once that such behaviour is compatible with the presence of ferroelectric domains in and out of plane on the same grain. As for the KPFM measurements, Fig. 3.3 - e) and f), the tip voltage has very large noise, making it difficult to draw any conclusions, but the second harmonic has a clear signal, showing the grains and their boundaries. The boundary between grains has a brighter contrast
more conductive than the grains themselves. This is plausible for a ferroelectric material. For sample B2, the grain size is tenfold greater (∼ 500 nm) than the other sample. The size difference can be attributed to the different substrate. The topographic images, Fig. 3.4 - a) and d), from KPFM and PFM are not of the same region since the process of measuring them requires the change of operation mode and this implies that the tip had to be retracted and approached once again selecting a slightly different area. This could be corrected for sample B1 with the tip piezo-actuators, but this was not possible for this sample. Despite that, there is an overlap between the two regions. The PFM images, Fig. 3.4 - b) and c), show that in-plane and out-of-plane polarization are different. The contrast of the out-of-plane component is more pronounced. Regarding the KPFM images, Fig. 3.4 - e) and f), they present stripes on the tip voltage and second harmonic. Artifact stripes are always along the scanning direction, in this case horizontal. The presence of vertical stripes could indicate that the sample is ferroelectric and the domain pattern is the one expected for an in-plane configuration.
Therefore, we can grow samples with thickness of approximately 500 nm that present ferroelectric domains both in-plane and out-of-plane. These results were important to further explore the impact that the substrate can have on the thin films, showing that the grain size can change depeding on the choice of substrate.
3.1.3
Group C: BaTiO
3on SrRuO
3/SrTiO
3The goal of this group was twofold: i) to optimize the electrodes for ferroelectric do-main measurements both in the atomic force and the photoemission electron microscopes; ii) check the sample growth reproducibility, fabricating two identical heterostructures (C1 and C2). The bottom electrode chosen was an approximately 40 nm thick layer of SrRuO3
Figure 3.4: PFM (red) and KPFM (blue) images of the sample grown on Nb:STO (B2). a) and d) represent the topography, b) is the piezoresponse signal phase and c) is the lateral piezoresponse signal phase. e) is the tip voltage and f) is the second harmonic. The grain size is approximately 500 nm.
Laser characteristics Deposition Annealing Name Sub. Fluency (J/cm2) # of shots T (◦C) T (◦C) Time (h)
C1 STO 1,3/1,2 10000/4000 770/756 655 2/1
C2 STO 1,3/1,2 10000/4000 770/755 655 2/1
C3 STO 1,3/1,2 2000/4000 770/755 655 0,5/1
Table 3.1: Parameters that changed during the samples growth. For the three bottom lines, the values separated by the bar are related to the BTO (to the left of the bar) and SRO (to the right of the bar) layer.
piezoresponse microscope tip. The samples were grown with the parameters shown in Ta-ble 3.1. Figure 3.5 shows the sample BS1 diffraction pattern. We can see the reflections from the BTO, SRO, and STO on the insets. The SRO presented only the (002) and (004) peaks, while BTO had all the (00l) peaks. Since the expected crystalline structure for SRO is pseudo-cubic just like BTO and STO, we believe that the (001) and (003) reflections are at a lower intensity, making it difficult to see in the diffractogram. The very thin peaks (glitches) that can be seem in the insets are due to an artifact of the
Figure 3.5: a) Diffraction pattern of the sample C1. The * show unidentified peaks. b) Inset showing the (002) reflections from the STO, BTO and SRO. c) Inset showing the (004) reflections from the STO, BTO and SRO.
Figures 3.6, 3.7 and 3.8 show the PFM images for the samples C1, C2, and C3, respectively. The first row is composed of the PFM images taken before the polarization and the second row by the images taken after. The process of polarization is described as follows: first, we used a voltage of 20 V between the tip and the bottom electrode to polarize a square of 25 µm2, then we used a voltage of -20 V to polarize a smaller square of
size 9 µm2 on the center of the first one. The polarization intensity can vary based on the
BTO layer thickness and the quality of the electrical contact between the electrodes and the microscope. All three samples showed remaining polarization, indicating that they are ferroelectric. The most intense polarization is the one in Figure 3.8, one explanation is that this sample has a thinner layer of BTO, approximately 20 nm. The other explanation for the differences between the polarization intensity is due to the contact between the eletrodes and the tip. This contact was made with indium wires and thus can be different for each measurement. The grain format is not much different for the samples BS1 and BS2. But sample BS3 has a unique format. And the surface roughness also changes from
Figure 3.6: Piezoresponse force microscopy images of the sample C1. a) and d) show the topography. b) and e), the phase of the out-of-plane piezoresponse and c) and f), the phase of the lateral piezoresponse. a), b) and c) were taken before trying to polarize the sample, while d), e) and f) were taken after polarization.
each sample, becoming smaller. That can be better seem when we look at the root mean squared (RMS) roughness values. These values are: 13 nm for sample C1, 10 nm for sample C2 and 5 nm for sample C3.
As shown by these results, we are able to grow ferroelectric BTO thin films with SRO bottom electrodes and gold titanium top electrodes. The ferroelectric domains of the samples could be patterned. Therefore, we have accomplished the first objective of the work, obtaining thin films with in-plane ferroelectric domains.
3.2
In-situ transfer
As described in section 1.4, we are interested to study the magnetoelectric coupling between a metallic Fe thin film grown on a BTO layer. In order to obtain such metallic
