Science, Faculty of
Physics and Astronomy, Department of
DSpace
UBCV
Roemer, Ryan Lee
Master of Science - MSc
University of British Columbia
Monolayer iron germanium telluride Fe₃GeTe₂, one of the typical two-dimensional ferromagnetic materials, hitherto, has only been studied by exfoliated micron-sized samples. We achieve high-quality wafer-scale growth of thin Fe₃GeTe₂ films by molecular beam epitaxy, greatly expanding the types of characterization tools employable and providing the possibility for its integration in devices like consumer electronics. Thickness-dependent transport measurements are used to characterize and probe for magnetic order. Ferromagnetic states exist in 1-10 layer thick Fe₃GeTe₂, with Curie temperatures ranging from 75 K in one layer samples to above 175 K in ten layer samples. A single ferromagnetic phase with signifi cant magnetic anisotropy is revealed for all layer numbers. We present the capability of synthesizing, wafer-scale Fe₃GeTe₂ as an essential step towards its fulfillment in any applications involving magnetism, such as spintronics.
Molecular Beam Epitaxy Growth andProperties of Fe3GeTe2 Thin FilmsbyRyan Lee RoemerB.Sc., The University of California, Santa Cruz, 2017A THESIS SUBMITTED IN PARTIAL FULFILLMENT OFTHE REQUIREMENTS FOR THE DEGREE OFMASTER OF SCIENCEinThe Faculty of Graduate and Postdoctoral Studies(Physics)THE UNIVERSITY OF BRITISH COLUMBIA(Vancouver)April 2021c© Ryan Lee Roemer 2020The following individuals certify that they have read, and recommend tothe Faculty of Graduate and Postdoctoral Studies for acceptance, the thesisentitled:Molecular Beam Epitaxy Growth and Properties of Fe3GeTe2 ThinFilmssubmitted by Ryan Lee Roemer in partial fulfillment of the requirementsfor the degree of Master of Science in Physics.Examining Committee:supervisor nameKe Zou, University of British Columbia (Physics and Astronomy)additional examiner nameAlannah Hallas, University of British Columbia (Physics and Astronomy)iiAbstractMonolayer iron germanium telluride Fe3GeTe2, one of the typical two-dimensionalferromagnetic materials, hitherto, has only been studied by exfoliated micron-sized samples. We achieve high-quality wafer-scale growth of thin Fe3GeTe2films by molecular beam epitaxy, greatly expanding the types of character-ization tools employable and providing the possibility for its integration indevices like consumer electronics. Thickness-dependent transport measure-ments are used to characterize and probe for magnetic order. Ferromagneticstates exist in 1–10 layer thick Fe3GeTe2, with Curie temperatures rangingfrom 75 K in one layer samples to above 175 K in ten layer samples. Asingle ferromagnetic phase with significant magnetic anisotropy is revealedfor all layer numbers. We present the capability of synthesizing, wafer-scaleFe3GeTe2 as an essential step towards its fulfillment in any applicationsinvolving magnetism, such as spintronics.iiiLay SummaryLay SummaryWhen layers of Fe3GeTe2 are grown in impurity free environments we areable to study the underlying physics of such a system free of the physicsassociated with defects. When these crystals are grown with precision it ispossible to control the number of layers size and shape leading to naturalintegration into devices. However, the growth of such crystals is no easytask. In this thesis we discuss the trials, failures and successes of the firstgrowth of ultra-thin films of Fe3GeTe2 revealing the intrinsic nature of thephase transition from the ferromagnetic to paramagnetic states.ivPrefaceThe work presented in this thesis was conducted during my tenure as amasters student researcher at the University of British Columbia under thesupervision of professor Ke Zou. It has involved collaboration with manymembers of this group. I specify this involvement, as well as other collab-orations and resulting publications which are detailed in this thesis. Muchof Ch. 1 and 4 is work conducted in UBC’s Thin Film Laboratory by my-self, Dr. C. Liu, and Professor Ke Zou. published as R.L. Roemer, C.Liu, K. Zou. Robust ferromagnetism in wafer-scale monolayer and multi-layer Fe3GeTe2, npj 2D Materials and Applications volume 4, 53 (2020).It details the growth and characterization of Fe3GeTe2 thin films grown bymolecular beam epitaxy.As the lead researcher on this project, I designed the project, grew thefilms, carried out measurements, and wrote the manuscript. Chong Liuprovided theoretical as well as experimental guidance and Ke Zou oversawand suggested the project. The two collaborators provided valuable feedbackduring the writing of the manuscript.vTable of ContentsAbstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . iiiPreface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . vTable of Contents . . . . . . . . . . . . . . . . . . . . . . . . . . . . viList of Tables . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . viiiList of Figures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . ixAcknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . xiiiDedication . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xiv1 Theoretical background of few quintuple layer Fe3GeTe2 12 Experimental Techniques & Design . . . . . . . . . . . . . . 52.1 Molecular Beam Epitaxy Technology . . . . . . . . . . . . . 52.2 Transport Measurements . . . . . . . . . . . . . . . . . . . . 92.3 Atomic Force Microscopy . . . . . . . . . . . . . . . . . . . . 142.4 Diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153 Growth Process; Exploring the Parameter Space . . . . . . 213.1 Sapphire FGT growth . . . . . . . . . . . . . . . . . . . . . . 263.2 Ge FGT growth . . . . . . . . . . . . . . . . . . . . . . . . . 333.3 Capping Decapping . . . . . . . . . . . . . . . . . . . . . . . 424 Experimental results . . . . . . . . . . . . . . . . . . . . . . . . 484.1 Structural Characterization of Pristine FGT Samples by Co-Deposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . 484.2 Transport results showing the conductivity and ferromag-netism in Fe3GeTe2 films . . . . . . . . . . . . . . . . . . . . 524.3 Temperature dependence of the anomalous Hall behavior . . 54viTable of Contents5 Discussion & Conclusion . . . . . . . . . . . . . . . . . . . . . 60Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62viiList of Tables2.1 Table of effusion Cells in both growth chambers in Zou’s Lab.Those above the double horizontal line live in Growth Module1 (GM1) those below are placed in Growth Module 2 (GM2). 7viiiList of Figures2.1 Examples of common shapes and sizes of MBE crucibles. Allcrucibles included above are made of pyrolytic boron nitride. 82.2 A non-exhaustive categorization of the conductive states yieldedby transport data . . . . . . . . . . . . . . . . . . . . . . . . . 102.3 A cartoon schematic of indium cold welded pad geometryplaced on the sample surface. . . . . . . . . . . . . . . . . . . 112.4 Post antisymmetrization the hysteresis loop shows signs ofthe normal Hall effect as indicated by the linear part of theloop at high magnetic field values. . . . . . . . . . . . . . . . 132.5 A typical hysteresis resulting from a transverse resistancemeasurement. . . . . . . . . . . . . . . . . . . . . . . . . . . . 142.6 An example of the experimental setup of an 2θ-ω scan. . . . . 162.7 Bragg scattering off a plane of atoms in a crystal. . . . . . . . 172.8 Common surface defects and their corresponding RHEEDdiffraction pattern signatures. . . . . . . . . . . . . . . . . . . 203.1 Lattice structure of Fe3GeTe2. . . . . . . . . . . . . . . . . . 223.2 The crystal structure of sapphire (Al2O3) . . . . . . . . . . . 233.3 AFM scan of the sapphire surface exemplifying islands left onthe substrate surface. . . . . . . . . . . . . . . . . . . . . . . . 243.4 Sapphire cleaned with methanol only. . . . . . . . . . . . . . 243.5 AMF height profile of the sapphire (0001) surface. . . . . . . 253.6 AFM scan of sapphire cleaned with acetone, isopropyl alcohol,and methanol. . . . . . . . . . . . . . . . . . . . . . . . . . . . 263.7 RHEED images of the sapphire surface along the two crystal-lographic directions where a) and c) belong to the [101¯0] direc-tion and b), d) belong to the [112¯0] direction. Top sub-figuresare pre-annealed substrates and the bottom sub-figures arepost-annealed RHEED images. Note: the RHEED intensitydramatically increases after annealing at high temperature. . 273.8 RHEED image of the first attempt at FGT growth. . . . . . . 29ixList of Figures3.9 Line profiles of sapphire substrates and possible films. . . . . 303.10 RHEED images of a polycrystalline film grown on the sap-phire (0001) surface. . . . . . . . . . . . . . . . . . . . . . . . 313.11 RHEED image of polycrystalline sample with 3D character. . 323.12 Examples of RHEED images with varying amounts of Te fluxin codeposition growth. . . . . . . . . . . . . . . . . . . . . . 333.13 The first truly 2 dimensional RHEED images obtained duringshutter growth on sapphire. . . . . . . . . . . . . . . . . . . . 343.14 RHEED images of Ge samples in the two unique crystallo-graphic directions. . . . . . . . . . . . . . . . . . . . . . . . . 363.15 2θ-ω scan from 0-100 deg. of the Ge(111) surface. . . . . . . . 373.16 RHEED images of FGT shutter method of 10 UC of FGTand other phases. . . . . . . . . . . . . . . . . . . . . . . . . . 373.17 XRD 2θ-ω scan from 0-90 degrees of shutter growth methodof a 10 UC film on Ge (111). . . . . . . . . . . . . . . . . . . 383.18 2x2 µm AFM image of a 10 UC FGT surface grown on Ge(111). 393.19 Transport data of FGT shutter method of 10 UC of FGT andother phases. . . . . . . . . . . . . . . . . . . . . . . . . . . . 403.20 RHEED images of FGT co-deposition method of 2 UC ofFGT and other phases. . . . . . . . . . . . . . . . . . . . . . . 403.21 2θ-ω scan of a 10 UC co-deposited FGT film on Ge(111). . . 413.22 2x2 µm AFM image of the FGT surface grown on Ge(111). . 423.23 Transport data of 10 UC FGT grown via codeposition withother phases. . . . . . . . . . . . . . . . . . . . . . . . . . . . 433.24 RHEED images of FGT film and Te cap. . . . . . . . . . . . . 443.25 RHEED image of film post Te decapping experiments. . . . . 453.26 RHEED images of a) 2 UC FGT sample and b) amorphousSe cap. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 453.27 RHEED image of a 2x1 reconstructed film surface during theSe decapping procedure. . . . . . . . . . . . . . . . . . . . . . 463.28 RHEED images of the Se decapped FGT films. . . . . . . . . 464.1 In-situ RHEED images of a 2x2 reconstructed Ge (111) sub-strates and FGT films . . . . . . . . . . . . . . . . . . . . . . 494.2 RHEED intensity oscillations during the growth of Fe3GeTe2 504.3 Ex-situ XRD of the out-of-plane lattice structure of FGT films. 514.4 Diffuse scattering FGT sample after exposure to ambient con-ditions for only a few minutes. . . . . . . . . . . . . . . . . . 524.5 A photo of an FGT sample measuring the longitudinal resis-tance with van der Pauw geometry in a PPMS . . . . . . . . . 53xList of Figures4.6 Longitudinal resistivity Rxx vs. temperature T taken of var-ious QL numbers. . . . . . . . . . . . . . . . . . . . . . . . . . 544.7 Fit of the Rxx with the Mott variable range hopping model,indicating the existence of localized states below 125 K. . . . 554.8 RAH vs. applied field of select QL numbers at 5 K showingthe hysterisis loops beloning to the AH measurement. . . . . 554.9 Anomalous Hall resistance of a single-QL, b 2-QL, c 5-QL,and d 10-QL FGT films. . . . . . . . . . . . . . . . . . . . . . 564.10 Transverse resistance measurement analysis. . . . . . . . . . . 58xiList of AcronymsList of Acronyms2D - 2-Dimensional3D - 3-DimensionalAH - Anomalous HallAFM - AntiferromagneticAHE - Anomalous Hall EffectARPES - Angle-Resolved Photoemission SpectroscopyDI - Deionized WaterFM - FerromagneticFGT - Fe3GeTe2GM1 - Growth Module 1GM2 - Growth Module 2HF - Hydrofluoric AcidIPA - Isopropyl AlcoholMAE - Magnetocrystalline Anisotropy EnergyMBE - Molecular Beam EpitaxyMRAM - Magnetic Random Access MemoryNC - Non-CappedPB - Parallel BeamPBN - Pyrolytic Boron NitridePG - Pyrolytic GraphitePMA - Perpendicular Magnetic AnisotropyPPMS - Physical Properties Measurement SystemQL - Quintuple LayerRAH - Anomalous Hall ResistanceRHEED - Reflection High Energy Electron DiffractionRMS - Root Mean SquareRR - Reminent ResistanceRS - Saturation ResistanceSC - Sticking CoefficientSTM - Scanning Tunneling MicroscopyUC - Unit CellUHV - Ultra High VacuumVdW - Van der WaalsXRD - x-Ray DiffractionxiiAcknowledgementsI would like to thank Dr. Chong Liu for his outstanding dedication to myscientific growth, Bruce Davidson, Fengmiao Li, Ted Shin and Simon Godinfor their thoughtful discussions and scientific insight. Finally I would liketo thank Professor Ke Zou for giving me a place to grow both literally andfiguratively.xiiiDedicationI dedicate this work to my mother, Melissa Robin Roemer, whose love andcare I will remember and carry through my life. Besos y abrazos.xivChapter 1Theoretical background offew quintuple layer Fe3GeTe2It is possible that the future of information storage and spintronics lieswithin 2-dimensional (2D) magnetic materials [1]. Spin transfer torques canchange the magnetization direction to control these systems. The ability tocontrol magnetic states electronically via spin transfer torques originates atthe interface. 2D materials pose potential for the increased density of inter-faces, as the surface-to-volume ratio becomes larger in reduced dimensions,making them desirable in applications. The ability to tune 2D materialsto display desired characteristics is also much greater. For example, strainscan be tuned in 2D, where 3-dimensional (3D) materials are of fixed bondlengths. This is important when constructing heterostructures where strainscan be used to increase interface compatibility. Another beneficial propertyof 2D materials is their stability in monolayer forms. The interlayer bondingis Van der Waals, thus the surfaces are free of dangling bonds. The samecannot be said of 3D materials where cleaving the largely ionic bonding be-tween layers often leaves the surface covered in dangling bonds. The abilityto tune magnetic order is also much greater in 2D materials as the mag-netic order is more robust in 3D materials, making manipulation of suchorders much more difficult. Magnetic metals, semiconductors and insulatorsall have prospects in applications. In magnetic insulators, long spin wavepropagation lengths are desired as well as the reduction in ohmic losses. 2Dmaterials also show strong proximity effects when coupled with 2D semicon-ductors yielding a way to control an extra degree of freedom:valley polariza-tion [2]. This degree of freedom has yet to be thoroughly explored outsideof a few materials.In the transition metal trihalides, particularly the Cr trihalides, all bulksamples, exhibit ferromagnetic order within layers with out-of-plane easy-axis directions for CrI3 and CrBr3 and in-plane for CrCl3 [3–5]. CrCl3 alsoexhibits antiferromagnetic interlayer magnetic order. The weak antiferro-magnetic coupling between the layers gives rise to extremely large tunnelmagnetoresistance [6].1Chapter 1. Theoretical background of few quintuple layer Fe3GeTe2In 3D, the integration of magnetic materials into electronic and photonicdevices has been taking place for decades [7–9]. One well-known exampleis the discovery of giant magnetoresistance and its applications [10]. Muchresearch has focused on engineering new magnetic structures, especially intwo-dimensions (2D). Prospects of which are plentiful in device application,not only in replacing magnets in current applications, but also in new struc-tures, such as in magnetic tunnel junctions and magnetic random accessmemories (MRAM) [11, 12].The discovery of graphene [13] inspired a new generation of materialsknown as Van der Waals (VdW) materials with intrinsic 2D nature [14].Samples were often accessible by exfoliating bulk crystals. After being pro-cessed as a device, the electronic properties of these 2D materials could becontrolled via strain, doping, gating, as well as “stacking” [15], to revealexotic states leading to an era of “VdW materials by design” [16]. Funda-mental physical properties have been realized that were proposed long-ago[17, 18], in 2D VdW materials, including the 2D superconductivity found inthin films of NbSe2 [19], FeSe [20], WTe2 [21], and Bi2Sr2CaCu2O8 [22]. Newopportunities in applications and a new generation of devices have becomeavailable due to the properties found in reduced dimensions.Exfoliated 2D flakes of CrI3 [23], Fe3GeTe2 [24], Cr2Ge2Te6 [25], andothers [26] have shown clear long-range magnetic order that are promisingfor applications involving properties such as above-room-temperature Curietemperature (Tc) and antiferromagnetic interlayer exchange.The library of such materials has recently grown drastically, and can beclassified as belonging to one of two categories; the transition metal halidesand the transition metal chacogenides. For example, the transition metaldihalides MX2 and transition metal trihalides MX3 have become extremelypopular as they have been shown to display a large range of magnetic be-haviours.Ferromagnetic (FM) order has been observed in both localized and itin-erant materials where all examples have out-of-plane magnetization. Exam-ples of which include the transition metal halides and chalcogenides [1]. Thekinds of Antiferromagnetic (AFM) order in 2D has become quite extensive.AFM zigzag, stripe, A-type and Neel AFM has been observed in variousstochiometries of transition metal halides and tertiary chalcogenides [1] Inthese materials, the transition metal (usually belonging to the 3D row) arearranged in a triangular lattice and bonded to nonmagnetic anions in thehalide column, namely iodine, chlorine and bromine. In these compounds,the halide atoms are stacked adjacent between layers in multiple configu-rations. The magnetic moments originating from the transition metal ion’s2Chapter 1. Theoretical background of few quintuple layer Fe3GeTe2spin and orbital angular momentum, interact via superexchanage where thefirst, second and third nearest neighbor superexchange interactions havebeen shown to be important in determining the kinds of magnetic order[27]. The superexchange is also quite sensitive to the cation anion bondangle and length.2D intrinsic magnetic order in a topological insulator has been observedin Manganese bismuth telluride (MnBi24Te4) [28]. The crystal structure iscomposed of a setuple layer structure which demonstrates in-plane A typeAFM order below 25 K [28] and FM order in-plane which undergo a spin-floptransition to a canted out-of-plane antiferromagnetic state.Yet, scalability and poor stability continue to hinder the progressiontoward the realization of the industrial-scale production of such devices [26].To date, the synthesis of many 2D magnetic materials still is all limitedto micron-sized flakes [29], eliminating any production scale applications.Here we focus on the wafer-scale synthesis of the VdW layered FM materialFe3GeTe2 (FGT), which has a Tc above room temperature [24] and is oneof the few 2D magnetic materials that shows evidence of topological states[30, 31].Single-crystal FGT was synthesized in 2006 [32], revealing a relativelyhigh Tc of ∼220K in bulk, in comparison to the absence of magnetism inthe isostructural Ni3GeTe2. This material has an out-of-plane magneti-zation vector with respect to the VdW planes in the ordered state [33],showing a large anisotropy defined as the energy dependence on the mag-netization vector direction. Calculations indicate its easy-axis is stabilizedby a large magnetocrystalline anisotropy caused by spin-orbit coupling [34].Angle-resolved photoemission spectroscopy (ARPES) measurements implythe FGT bulk belongs to a class of topological nodal line semimetals [35].Another ARPES experiment suggests the FM state is accompanied by aheavy-fermion state, which has only been seen in 4 f and beyond electronsystems [36].Exfoliated FGT monolayers and multilayers [24, 37] preserve the char-acteristic FM order of bulk. Tc reaches 310K upon ionic liquid gating intrilayer samples [24], one of the highest FM transitions recorded amongknown 2D materials. Unlike CrI3, no interlayer exchange coupling hasbeen observed in FGT films. It has been shown that FGT hosts a largeBerry curvature, which is responsible for the anomalous Hall effect (AHE)[35]. Thin-film measurements substantiate claims of topological propertiesin the observation of Ne´el-type skyrmions in FGT/WTe2 heterostructuresand large anomalous Nernst effect [31, 38].Until recently, FGT films remained the only metallic, rare-earth free, and3Chapter 1. Theoretical background of few quintuple layer Fe3GeTe2layered FM with out-of-plane magnetization in 2D. Because of the metal-lic characteristics, its fundamental magnetic properties can be studied viatransport measurements, together with optical measurements such as themagneto-optic Kerr effect. Metallicity also enables easy control of charge aswell as spin degrees of freedom by electrical currents, useful in device ap-plications. For example, FGT films have been made into ultrathin devices,realizing current control of the perpendicular magnetization via spin-orbittorque, with efficiencies comparable to that of 3D FM heterostructures-thecurrent leader of the field [39]. This makes FGT all the more desirablefor magnetic random access memory (MRAM) [40]. Closely related appli-cation prospects are found in magnetic tunnel junction heterostructures ofFGT/graphite/FGT, which display promising characteristics of 160% tun-neling magnetoresistance at low temperature [41].We employ the molecular beam epitaxial technique to synthesize atom-ically smooth and single crystalline FGT films, with thickness ranging fromone quintuple layer (a single QL, also the smallest substructure of FGT,∼ 0.8 nm) to ten QLs on 5 x 5 mm substrates. Among reported monolayersamples, ours display the highest Tc probed by AH measurements, togetherwith an order of magnitude larger coercivity [24, 42]. Large switching ratioshave been found in all samples at low temperatures as well as significantmagnetocrystalline anisotropy. Our wafer-scale samples yield the opportu-nity for measurements and applications, which can be extremely difficult onthe micron-sized flakes yielded by exfoliation techniques, such as ARPESand x-ray spectroscopies. One can begin integrating FGT films into com-mercial devices following the established techniques of silicon.4Chapter 2Experimental Techniques &Design2.1 Molecular Beam Epitaxy TechnologyUltra High VacuumIn this work, we focus on the growth of 2D materials using state-of-the-arttechnology known as, molecular beam epitaxy (MBE). MBE takes place instainless steel chambers which are kept in high or ultra high vacuum (UHV)environments which we define as pressures ranging from 10−3 to 10−8 torr,and less than 10−8 respectively. These pressures result in gas densities of theorder 1013 and 107 mol/cm3 respectively. To contrast this point the densityof molecules in abient conditions is roughly 2x1021 mol/cm3 In UHV fewspecies of atoms remain, the majority of which are CO N2 and H2. Atpressures so low the mean free path of atoms and molecules, which survivein the chamber, approaches the tens of miles. These densities allow forgrowth of the most pristine, impurity-free films accessible to date which arerequired not only to study and disentangle fundamental physics from thatof disorder and impurities, but are also necessary for the semiconductingindustry today.To achieve such low pressures, specialized vacuum pumps are required.To approach UHV, it is common to use a roughing and turbo pump in se-quence and often others such as cryopumps and ionpumps are necessary.Even then, specialized pumps may be required to further vacate the cham-ber.Sources, Source Materials and Evaporation TechniquesIn MBE, pure source materials are epitaxially transported to the samplesurface in the gaseous state. The technique used to sublimate the sourcesvaries, and depends on parameters such as the vapor pressure of an elementin the liquid and solid state. Some elements require sub-zero temperatures52.1. Molecular Beam Epitaxy Technologyto maintain the solid state, while others are required to be heated past 1000◦C before the vapor pressure yields a measurable atomic flux. In general,elements with d and f shell valence states have higher melting points wheres and p shell elements tend to have lower melting points (with the excep-tion of some mono- and divalent p series elements). Fortunately, a varietyof techniques have been developed to satisfy such requirements. For themajority of elements, a crucible is sufficient in holding the source materials.In Zou’s lab, three volumes of crucible are used, 5 cc, 10 cc and 30 cc. Acommon but non-exhaustive list of such materials which crucibles are madeinclude alumina, graphite, pyrolytic boron (PBN), pyrolytic graphite (PG),tantalum, titanium, and tungsten based on the material placed inside. Forour purposes, we use conical crucibles with exceptions for Se. These cru-cibles are wrapped with resistive heaters and monitored by therocouples toachieve atomic sublimation and maintain consistent flux. The combinationof heaters crucibles and thermocouples are often referred to as “cells”.Cells are further characterized by the number of filaments and types ofelemental metals used in their construction. High temperature-low vaporpressures source materials require refractive metals to operate and are re-ferred to as high temperature cells. Single filament cells or low temperaturecells operate between 100◦C and 750◦C and have a small thermal mass toensure stability. This is particularly important when dealing with high va-por pressure materials because slight variations in thermal energy can resultin relatively large changes in atomic flux.Dual filament cells allow one to regulate different zones of the crucible.For example, during the growth of aluminum containing compounds, thelip of the crucible is kept cooler then the base. This ensures the sourcematerials remain inside the crucible and refrain from wetting any otherMBE components. In Zou’s lab we strictly use single filament cells.Single fillament high temperature cells operate up to 2000◦C while ensur-ing constant temperature stability. This is enabled by high purity tungstenand tantalum hot zone at the base of the crucible.For specialty elements such as chalcogenides, a cracker is used. Currentlya 500 cc valved cracker is used in the Zou lab for Se in the growth of Secontaining films. Selenium will vaporize as Se2 similar to O2, as they bothhave the same valence state. Thus, after the initial sublimation from bulkSe to diatomic Se, the atomic flux is carried though a region where the wallsmay be maintained at a higher temperature relative to the bulk state, suchthat when diatomic selenium undergoes a collision with the walls, it “cracks’in two and becomes atomic selenium, which is then epitaxially delivered tothe sample surface. The cracker also is equipped with a needle valve such62.1. Molecular Beam Epitaxy TechnologyMaterial Cell Type CruicibleSr 30cc Low Temp PBNBi 30 cc Standard PBNNd 10cc High Temp WLa 10 cc High Temp WNi 10 CC High Temp W + BeOEu 30 cc Standard PBNTi 10 cc High Temp WTe 30cc Low Temp PBNGe 30cc Standard PBNFe 10 cc High Temp W + Al2O3Table 2.1: Table of effusion Cells in both growth chambers in Zou’s Lab.Those above the double horizontal line live in Growth Module 1 (GM1)those below are placed in Growth Module 2 (GM2).that the high vapor pressure material may be completely isolated from thegrowth chamber and an atomic flux can be more easily regulated.For convenient reference the table 2.1 reports the current state of thesources and crucibles in our dual GENxplor system.Veeco manufactures other crucible types but for our purposes these op-tions suffice. These crucibles are further customizable in the material as toadhere to each materials specific needs. Elements with high melting pointssuch as tungsten W (∼ 3000◦C) require special procedures to sublimate thematerial like an electron beam evaporator.SubstratesIn the MBE chamber, materials are deposited onto another crystal knownas the substrate. This crystal, as discussed below, is a key choice in thegrowth process. Most importantly, it provides a template to grow. We canbreak the substrate and film relationship into two categories: epitaxial andnon-epitaxial growth. Non-epitaxial growth is a growth where the film andsubstrate have little to no coherence. In epitaxial growth the substrate ischosen such that the film grows in the orientation of the substrate. This isdone by choosing a substrate with similar in-plane lattice constants. The filmthen adheres to the substrate such that each unit cell falls onto some integermultiple of the unit cells of the substrate. In this work many substrates are72.1. Molecular Beam Epitaxy TechnologyFigure 2.1: Examples of common shapes and sizes of MBE crucibles. Allcrucibles included above are made of pyrolytic boron nitride. Figure adoptedfrom [43]used.Strain & Interface EngineeringThis template can be chosen such that the film and substrate have similarlattice constants which we then expect the film to grow epitaxially on thesurface. Although the film growth may be epitaxial, minute differences inlattice constants will lead to in-plane strain on the film.In this case, substrates not only provide a template to grow the targetmaterial, but can often lead to new physics due to a variety of reasons.The substrate can be chosen to either stretch or compress the film. Thisallows one to explore physics that may be emergent given a proper strain.Oftentimes this leads to interesting phenomena in otherwise banal system.These strains are broken into two categories: compressive and tensile strain.Compressive strain indicates the films in-plane lattice constants are slightlysmaller than those of the chosen substrate. Tensile strain implies the op-posite. The induced strain can often lead to new behaviors or changes inphysical properties of the film. As the film becomes thicker and more unitcells are deposited, the strain begins to relax through a variety of mecha-nisms, eventually relaxing to its bulk values.82.2. Transport MeasurementsPreparation to Ensure Atomically Clean SurfacesOftentimes substrates arrive from distributors with dirty surfaces. Dirtyhere meaning atomically rough or littered with contaminants such as organ-ics on the surface. To rid the surfaces of such contaminants various etchingprocedures are used. These procedures often require the use of extremelyacidic solutions like HCl or H2SO4, but many contaminants can be removedwith volatile chemicals such as acetone and ethanol.To make the substrates atomically smooth high temperature annealingprocesses may be required either in UHV or with flows of ultra-pure gassessuch as O2 or HAr.2.2 Transport MeasurementsAll transport measurements were carried out in a Quantum Design physicalproperties measurement system (PPMS), where Van der Pauw geometrywas used to measure longitudinal (Rxx) and transverse (Rxy) resistances.Contacts were made via indium cold pressing and attached via a gold wire.Van der PauwThe Van der Pauw geometry is a common technique to probe the resistanceof approximately 2D samples. In this limit, the technique allows one to probesamples of arbitrary shape granted the film is continuous with a negligibleamount of holes. In this design, four contacts are placed on the perimeter ofthe film. Although this approach does not measure resistivity of a specifiedarea, it yields a resistance averaged over the entirety of the film. This maybe a problem for anisotropic materials, but in MBE atomic, uniformity isassumed over the entire growth area. The four contact placement allows formultiple types of measurements from which properties may be extracted.Such properties and the techniques to extract them are detailed below.Longitudinal GeometryResistances may be measured in a number of different geometries and eachgeometry probes for a unique property. Longitudinal resistance quantifiesthe insulating, conducting or semiconducting properties of a material. Themeasurement is often performed as a function of temperature where a con-tinuous current is supplied between two contacts, while the remaining pairof contacts measures the potential difference on the same surface as shown92.2. Transport Measurementsby Figure 2.3 a). This allows one to decouple the contact resistance fromthe intrinsic resistance of the sample. As the current is passed through thesample, voltage is read at each temperature to determine the resistive be-havior of the sample. The resistance of a sample may be used to categorizeit in term of its resistivity as shown in Figure 2.2Figure 2.2: A non-exhaustive categorization of the conductive states yieldedby transport data adopted from [44].Sheet resistanceMany thin films, especially Van der Waal’s materials, show conductivitywithin sheets. One helpful way to analyze materials is by the resistance weexpect to be in a 2D planes. For example, in a 2D electron gas, it is notobvious what the thickness of the conducting plane is, thus we analyze theresistance in terms of sheet resistance. The units of such are ohms. ManyVdW materials are quasi 2D electron gasses. Thus sheet resistance can behelpful in extracting properties such as carrier density and mobility.R =ρtLW= RSLW: RS =ρt(2.1)102.2. Transport MeasurementsFigure 2.3: a) A cartoon schematic of indium cold welded pad geometryplaced on the sample surface. Note the V+, V- and I+, I- groups lie inparallel, indicating a longitudinal geometry. V and I stand for current andvoltage respectively. b) A cartoon schematic of transverse pad geometryplaced on the sample surface. Note the V+, V- and I+, I- groups areorthogonal indicating a transverse geometry. V and I stand for current andvoltage respectively.Here R is the measured resistance, ρ is the resistivity, L, and W are thelength and width of the sample respectively and RS is the sheet resistance.Transverse geometryHall ResistanceAnother useful quantity that is commonly extracted from transport mea-surements in the Van der Pauw geometery is the Hall resistance. The Hallresistance yields valuable information such as the type of carrier (i.e. elec-trons or holes) as well as the concentrations of such carriers. The Hallresistance is a result of the Hall effect, in which a current subject to anapplied magnetic field leads to the eventual buildup of carriers on one sideof the sample due to the Lorentz force detailed in Figure 2.2F = q(E+ v ×B) (2.2)The direction of the resulting voltage is determined by the carriers in thesample. By fitting the linear portion of the transverse resistance in theanamalous Hall measurement detailed below, it is possible to extract theHall coefficient. It is possible to describe the Hall resistance in terms of the112.2. Transport Measurementscurrent density and in turn the carrier density of single carrier materials asfollows. Let us define a sample subject to out-of-plane magnetic B field inthe positive z direction, while current I is carried in the positive x direction.RH =EyjxB=1ne(2.3)Anomalous Hall MeasurementsThin films can be notoriously difficult to measure using standard magnetom-etry, this is due to the relatively small number of magnetic atoms relativeto bulk.In metallic materials, such as Fe3GeTe2, this challenge can be circum-vented with transport measurements. To do this, one places contacts inthe transverse Van der Pauw geometry, identical to a Hall measurement asshown in Figure 2.3 b). Samples are then cooled and may experience astructural magnetic or electronic phase transition. Once a desired temper-ature is reached, the samples are then isothermally exposed to a magneticfield while current is ran through the sample. The applied field is sweptthrough a desired cycle (usually a few T). In an intrinsic ferromagnet thismeasurement theoretically results in a transverse resistance signal,Rxy = RN (H) +RAH(M) (2.4)where Rxy is the measured signal from the PPMS (H)is the applied fieldRN is the normal Hall resistance(H) is the applied field. Magnetization isthen related to the second term where RAH is the AH coefficient and (M)is the magnetic signal from the sampleThe AHE occurs in solids where spin-orbit coupling exists, which breaksthe time-reversal symmetry. Multiple mechanisms are theorized to play arole, including Berry-phase related physics and disorder related phenomena.Berry phase AHE are said to dominate in metallic ferromagnets with lowconductivity which is the case for Fe3GeTe2.Raw data yielded by experiment are usually not antisymmetric as onewould expect from a non-interacting magnetic hysteresis. These shapes aredue to imprecise pad geometry, in other words, a typical Rxy measurementcontains Rxx components. Thus raw data were first antisymmetrized toremove the Rxx component, according to122.2. Transport MeasurementsRH↑(B) = [Rxy↑(+B)−Rxy↓(−B)]/2 (2.5)andRH↓(B) = [Rxy↓(+B)−Rxy↑(−B)]/2(2.6)Here Rxy,↑ and Rxy,↓ are the values during the increasing and decreasingfield sweep, respectively, and B is the applied magnetic field.After such an antisymmetrization, the loops saturated resistances in thenegative and positive field directions have equal magnitudes, as displayed inFigure 2.4. Here it is important to note the loops are centered about theorigin, but there remains a linear portion with a non-zero slope where theloops have become saturated in the high field regions.Figure 2.4: Post antisymmetrization the hysteresis loop shows signs ofthe normal Hall effect as indicated by the linear part of the loop at highmagnetic field values.This linear portion slope shown in Figure 2.4 is due to the normal Halleffect from which the Hall coefficient may be extracted to reveal the carriertype and concentration. In anamalous Hall experiments this contributionis removed. In this experiment the normal Hall coefficients were subtracted132.3. Atomic Force MicroscopyFigure 2.5: A typical hysteresis resulting from a transverse resistance mea-surement. Note: the non-zero slope of the saturated linear portion of thehysteresis and the non-symmetric position of the loop due to the Rxx contri-bution and non-exact pad placement. The yellow circle indicates the valueof the remanant resistance, the red triangle indicates the value of saturationresistance and the blue square is the coercive fieldby fitting the magnetically saturated portion of the data with a term pro-portional to the applied field. What remains is an anomalous resistance,which is directly related to the intrinsic magnetization, an example of whichis plotted in Figure 2.5. Here is an example of a magnetic hysteresis loopobtained from anamalous Hall measurements on a film in the FM state aftersymmetrization and decoupling tricks have been performed. Labeled in thefigure are a few quantities that may be noted. The yellow circle representsthe remanant resistance, the red triangle marks the saturation resistanceand the blue square is the coercive field value.2.3 Atomic Force MicroscopyAtomic force microscopy (AFM) is a kind of scanning probe microscopy ca-pable of imaging sample surfaces at the nanometer scale. In our experiments,we used a tapping-mode technique to avoid contamination and destructionof the sample surface. The apparatus operates using a silicon tip usuallymicrometers in diameter, placed at the end of a cantilever, which then oscil-lates vertically (z direction). The oscillations of the cantilever reflect a laserto a detector as the tip scans in the x-y plane. Deflections from these oscil-142.4. Diffractionlations due to electrostatic forces between the tip and sample are recordedand used to map the sample surface via an image by assembling lines ofhorizontal scans.In our experiment, AFM was helpful to investigate the surface topogra-phy of substrates to ensure atomic smoothness in hopes of obtaining atom-ically flat films of such substrates. It was also beneficial in determining theeffects of pre-growth etching and annealing procedures.2.4 DiffractionIn the Zou lab, samples are probed with x-ray diffraction (XRD) ex-situ.From X-ray diffraction, one is capable of determining the film thickness,crystal structure and in the case of heterostructure growth, the superlatticeperiod. In this experiment, high energy photons are emitted from a Cuanode. The Cu anode is the target of the electrons which knock out coreelectrons. Outer electrons drop down to fill the empty orbital and emitx-rays. The photons emitted are not monochromatic. In our case, bothkα1 and kα2 emission channels are used, with wavelengths of 1.5406A˚ and1.544425A˚ respectively. According to our diffraction conditions these twowavelengths result in a superposition of two diffraction peaks. In theory,this results in a tightly spaced double peak. In practice, these peaks areextremely difficult to resolve at low angles, thus the average wavelength ofthe two peaks is used to determine lattice constants. Some diffractometersare equipped with a monochromator which suppresses the kα2 emissionchannel. The trade off here being the total intensity is diminished, but theprobing wavelength is exact, allowing for a more precise determination ofthe lattice parameters. This minute difference in lattice parameters canbe problematic as some imprecision caused by the mixed wavelengths canbe of similar order of magnitude as the question at hand in cases such asdetermining relaxations, and differences in bulk and thin film structures.Once emitted from the anode, the photon is diffracted by the sample.The incident beam penetrates a few microns due to the energy of the pho-tons. This is why substrates often generate the largest peak intensities(many orders of magnitude larger than the film). The majority of thediffracted beam penetrates a few microns below the surface, a great dealdeeper than the usual thicknesses of films. To further highlight this point,signals from the metal sample holder can also be observed. This is whysamples are placed on an anti-reflective Ge(111) zero-diffraction disk whichprevents secondary diffraction processes. Once diffracted, the photons travel152.4. Diffractionto a detector.For the purposes of this thesis 2θ-ω scans were performed. In this tech-nique, both the detector and source are coupled and move symmetricallythroughout the scan as displayed in Figure 2.6. Note: ω is not necessarilyequal to 12 ∗ 2θ because there is an intrinsic offset of the sample. This offsetoriginates from the substrate having its diffraction vector (marked S in Fig-ure 2.6)offset with respect to the normal of the sample holder. This couldbe due to miscuts of the substrate or residual paste on the backside.Figure 2.6: An example of the experimental setup of an 2θ-ω scan. Omegais defined as the angle of the beam with respect to the horizontal while 2θ isthe angle of the incoming beam with respect to the detector. The diffractionvector marked S.In 1913 Bragg & Bragg attributed the characteristic patterns of reflectedX-rays to parallel planes of atoms, which composed the crystal. They hy-pothesized that such planes were equally spaced, and consequently, atomsreflecting x-rays would do so constructively and destructively based on theangle of incidence. For monochromatic light, constructive interference isonly possible when an integral number of wavelengths fit in the extra dis-tance a wave reflects off successive planes. This has become known as theBragg condition, as shown in eq.2.7nλ = 2dsinθ (2.7)Here n is an integer, λ is the wavelength of the incident photon, and θ isthe angle of incidence measured from the lattice plane, as shown in Figure2.7.For the purposes of this thesis, XRD was used to determine out-of-planelattice constants ex-situ.162.4. DiffractionFigure 2.7: Bragg scattering off a plane of atoms in a crystal. Note: theextra distance traveled by the incident wave interacting with the lower planeis 2dsinθ figure adopted from [45].RHEEDWhen growing materials with multiple elements, there may exist compoundswith similar in-plane lattice constants and symmetries with different stoi-chiometries. Thus during growth, it is impossible to be determine what hasbeen grown in-situ with 100% confidence. The uncertainty can be greatlydiminished by measuring the in-plane lattice constants by diffraction.During the growth process, a powerful tool to monitor the sample condi-tion is reflection high energy electron diffraction (RHEED). It allows users tooptimize stoichiometry as well as determine surface termination for layeredmaterials. Contrary to XRD, this technique uses electrons emitted from anelectron gun to diffract from the first few layers of a sample surface at graz-ing incidence (in some systems this process can be quite destructive). Thediffracted electrons are then accelerated to a phosphorescent screen where acharacteristic diffraction pattern is revealed. Thus the process is electron in,electron out. The diffraction patterns are a result of the wave-like nature ofthe electrons interacting with the periodic electric potential created by thearray of atoms. To probe the crystal for its atomic structure the electronswavelength must be comparable to that of the lattice spacing which is of theorder of a few angstroms. The relationship between a particles momentumand wavelength is given by the de Broglie wavelength λ = hmv where h isthe plank constant h = 6.626 E 34 m2kg/s, m is the electron mass 9.109 E31 kg which are both constants. Thus to control the de Broglie wavelength,the only free parameter we have to tune is the electron velocity, which is172.4. Diffractioncontrolled by the potential difference supplied by the RHEED apparatus.In our experiment, we use an energy of 15 KeV which results in an electronwavelength of 0.122 A˚. which is an order of magnitude less than most latticespacing. With a sufficient potential difference supplied to accelerate elec-trons the next step is to generate enough electrons such that the patternsdisplayed on the phosphorescent screen is of adequate intensity. This is donevia control of the emission current. In our experiment, we use an emissioncurrent of roughly 0.3µA, which is sufficient to characterize the diffractionpattern while avoiding the potential pixel burnout on our screen. In somesamples containing highly volatile elements (those with high vapor pressurelike Se) it is not uncommon to create vacancies in the material or completelydestroy the lattice altogether. This must be kept in mind while adjustingthe RHEED intensity and energy. Generating diffraction images withoutdamaging samples has been eased by the data acquisition software in KSA400 system. The program allows users to adjust the exposure time per frameas well as sum over multiple frames. The trade-off with this approach is, thegrowth process becomes less interactive because images are generated everyfew seconds rather than multiple images per second.Often times the surface of a film will not resemble its bulk structure.Atoms on the surface will rearrange themselves to find a lower energy statewhich is commonly referred to as surface reconstruction. Fundamentallythe rearrangement results from a discontinuity in the potential above thesurface layer. Thus the anharmonic forces which determine the equilibriumpositions move the atoms in either the lateral or vertical direction referred toas relaxation or reconstruction respectively. Relaxation among most metalsbrings the surface atoms closer to the existing crystal. In other words, thec lattice constant will slightly decrease for the surface layer. Intuitivelythis makes sense because the interlayer bonding below is not competingwith a layer above. For a certain crystal structure there may be multiplereconstructions that emerge under various conditions (i.e. temperature).These reconstructions hint at long range order within the beam spot. Theprototypical example of which is the 7x7 Si(111) reconstitution.Because each crystal structure has a unique diffraction pattern, it is pos-sible to extract various properties of the crystal under consideration. Thisis particularly important when growing crystals with MBE. For example, el-ements or compounds evaporated from crucibles may form multiple crystalstructures when combined, sometimes with the same atomic flux. Luckily,RHEED is capable of distinguishing between phases with its multitude oftools. In many cases, the phases may be identified in-situ by their uniquecrystal structures and lattice parameters allowing for a more interactive182.4. Diffractiongrowth. Unfortunately, due to the grazing incidence, the c lattice parame-ter is inaccessible, yet in-plane lattice constants are accessible by analyzingthe distances between diffraction spots. This approach is best done by com-paring the known in-plane lattice constants of the bulk crystal substrate. Inthe KSA software there exists a pixel spacing tool in which users can placea line with variable width (to integrate and average) to quantitatively studythe in-plane lattice spacing.Upon growth of our 2D materials starting from a bulk substrate, it ispossible to witness the diffraction pattern change in real time. This is be-cause the underlying reciprocal space rods from the top layers reveal arereplaced with those corresponding to the new material.This situation presents the opportunity to analyze the as grown in-planelattice constants. Because commercially available bulk substrates are usu-ally of high quality, it is typical within the community to accept the welldocumented bulk lattice constants. This paired with a constant sample ge-ometry allows one to trace the progression of lattice constants by simplynoting the ratio of pixel spacing and known bulk constants. This providesan immediate check as to the in-plane constants.Often times, coherently grown epitaxial films will naturally be straineddue to lattice mismatch at the substrate/film interface. This strain willbegin to relax after just a few layers and will eventually relax to its bulklattice constants. This strain evolution can sometimes vary as much as 10%of the bulk lattice constants.Surface defects can greatly influence the diffraction pattern. Qualita-tively diffraction patterns may be characterized by the eye, as described inFigure 2.8 where common surface defects are listed with the correspondingchanges to an otherwise ideal diffraction pattern in Figure 2.8.The other important information such as the crystal symmetry can befound by RHEED. This is done by noting the number of unique diffrac-tion patterns found after rotating the sample a complete revolution. Sam-ple quality can often be inferred by the quality of the diffraction pattern.Streaky lines indicate 2D growth where point like diffraction spots indicateatomically smooth films.192.4. DiffractionFigure 2.8: Common surface defects and their corresponding RHEEDdiffraction pattern signatures figure adopted from [46] .20Chapter 3Growth Process; Exploringthe Parameter SpaceIn this chapter the unsuccessful growths of Fe3GeTe2 FGT will be discussed.Various parameters such as substrate temperature, atomic fluxes and growthperiods will be explored. These trials provided valuable insight into thegrowth process and how to efficiently explore the parameter space. Thischapter is provided for the reader such that future growths of similar com-pounds will be easier upon comparison to data presented. Conditions thatultimately led to the successful growth of FGT are discussed in chapter 4.FGT films were synthesized using a Veeco GENxplor system. Standardeffusion cells are mounted at the bottom of the chamber facing the samepoint at the center of the chamber where the substrate is facing downwards.Source materials of Fe (99.98%), Ge (99.999%), and Te (99.9999%) wereevaporated from the effusion cells.The FGT unit cell is described by a hexagonal symmetry (space groupP63/mmc) composed by two weakly coupled quintuple layer (QL) substruc-tures, with bulk lattice parameters a = b = 3.99 A˚, and c = 16.33 A˚, asshown in Figure 3.1. The middle three layers of each QL form a heterometal-lic slab of Fe2+Ge4− surrounded by Fe3+ atoms. Encapsulating the slabs aretellurium layers. Theoretical calculations of monolayer magnetic momentsyield Fe3+ and Fe2+ atoms contribute 1.723 and 1.005 µB, respectively [34],where is the Bohr magneton. The structure is displayed in Figure 3.121Chapter 3. Growth Process; Exploring the Parameter SpaceFigure 3.1: Lattice structure of Fe3GeTe2 (FGT), where a single unit cell(two quintuple layers (QL)) is outlined. a) top down view of the FGTunit cell and right: the lattice projected along the a axis . There existtwo inequivalent Fe sites, Those closest to the tellurium layers have a 3+oxidation state where those in the Ge layer have an oxidation state of 2+.The 3+ state are said to carry larger magnetic moment with respect to the2+ state [34].In early trials of the growth of FGT, Al2O3 (sapphire) was selected as thesubstrate of choice following suit of the only reports of MBE grown FGT asin reference [47]. Here the reader should note the alternating layer structureof sapphire and the inter layer distance of roughly 2 A˚. The commerciallybought substrates were of (0001) termination. Yet, due to miscuts incurredduring fabrication, the termination layer alternates between the A and Blayers indicated in Figure 3.2. Although the in-plane symmetry of sapphireand FGT are both hexagonal, the in-plane lattice parameters of sapphireare a=b=4.758A˚ roughly 18% larger than that of FGT. From this large dis-crepancy it is safe to assume the film would not be coherently strained. Yet,sapphire is often a good substrate choice because it can be easily prepared toatomically smooth terraces and non-polar surface. Substrates were orderedfrom MTI Corp as well as Crystec. The differences between the two sub-strate supplied were mostly negligible, the only comment-able difference wasthe reported c axis parameter. From MTI the reported c lattice constant was22Chapter 3. Growth Process; Exploring the Parameter Space12.992 A˚, while the Crystec c parameter was given to be 13.002A˚. Fortu-nately this difference would not pose any problem. As both could eventuallyyield atomically flat surfaces. The original procedure used an initial washFigure 3.2: The crystal structure of sapphire (Al2O3) The unit cell is com-posed of 8 alternating A and B layers. The Lattice constants are a=b=4.763A˚, and c=13.003 A˚. Figure adopted from [48]in acetone followed by an isopropyl alcohol (IPA) wash in a sonicator for 3minutes. The results left the surface rough as displayed in the AFM scan indisplayed in Figure 3.3. A methanol wash was also performed and comparedto the acetone and IPA. The results were slightly more desirable with lessislands in a larger scan size as displayed in Figure 3.4. The resulting surfacesdid not display atomic terraces conducive to growth.23Chapter 3. Growth Process; Exploring the Parameter SpaceFigure 3.3: A 2x2 µm AFM scan of the sapphire surface exemplifyingislands left on the substrate surface. a) a forward height scan of the sapphiresurface where islands are clearly visible as indicated by the round spots ofbrighter hue. b) the tapping phase in the forward direction which serves tohighlight the non-uniformity of the surface.Figure 3.4: Sapphire cleaned with methanol only. a) A 2x2 µm scan ofthe Sapphire surface.b) The height profile of the diagonal white line in a)the height profile ranges roughly 1000 pm over the entire scan. The RMSsurface roughness is roughly 870 pm.Because the surface did not display atomic terraces a high temperatureannealing process was developed to reveal atomic terraces. Substrates wereplaced in a furnace with an oxygen flow at 1200 ◦C for 30 minutes. The24Chapter 3. Growth Process; Exploring the Parameter Spacewarm-up was linear at 6 ◦C/min and the cool-down was also linear at 4 ◦C/minute which resulted in nice atomic terraces, as displayed in Figure 3.5 inwhich the sapphire crystal demonstrates atomic terraces over the entiretyof the scanned surface. The technique has further been developed using amethanol rise in the sonicator for 3 minutes repeated 3 times subsequentlyfollowed by the annealing procedures described above with correspondingscans displayed in Figure 3.6. The result was large atomically smooth ter-races with no apparent surface decoration i.e. no islands. With the resultingsubstrates it was possible to begin the growth process.Figure 3.5: A 2x2µm AFM height profile of the sapphire (0001) surface.a)Clearly resolved atomic terraces. This sample was placed in the furnaceat 1200C for 30 minutes. The heat-up and cool-down were over a 110 and600 minute duration respectively. The image to the left is the height wheredarker represents low height and bright areas represent high portion. b)Theheight profile along the white line indicated in figure a) . The steps alongthe line correspond to a single atomic layer in the sapphire crystal along thec axis.253.1. Sapphire FGT growthFigure 3.6: AFM scan of sapphire cleaned with acetone, isopropyl alcohol,and methanol followed by a high temperature annealing at 1120circC for 60minutes. More specifically a 110 min heat-up, and 600 min cool-down. Theheight variation of the scan is roughly 6 A˚.3.1 Sapphire FGT growthAfter the surfaces were cleaned, the substrates were ready to be degassedin UHV. Substrates were mounted on either tungsten or molybdenum cir-cular sample holders with carbon paste that used water as a solvent andloaded into the loadlock chamber. In the loadlock chamber, substrates wereheated by two argon gas filaments to 200◦C for 2 hours. Prior to degassing,pressures were roughly 10−7torr and after degassing would reach ∼10−9torr. After the initial degassing procedure was completed, substrates werebrought through the buffer chamber and into GM2, the chalcogenide growthchamber. In GM2, samples were further degassed at 600 ◦C for 30 minutes.RHEED images were then taken of the two distinct crystallographic direc-tions. The [101¯0] direction as well as the [112¯0] directions both pre-growthand post-growth as displayed in Figure 3.7. Here a) and b) correspond to263.1. Sapphire FGT growthpre-annealed substrates and c) and d) correspond to substrates annealed at600◦C for 30 minutes. a) and c) correspond to the [101¯0] crystallographicdirections, while b) and d) correspond to the [112¯0] direction. All RHEEDimages were taken with a 0.3 A emission current.Figure 3.7: RHEED images of the sapphire surface along the two crystallo-graphic directions where a) and c) belong to the [101¯0] direction and b), d)belong to the [112¯0] direction. Top sub-figures are pre-annealed substratesand the bottom sub-figures are post-annealed RHEED images. Note: theRHEED intensity dramatically increases after annealing at high tempera-ture. .Codeposition vs Shutter GrowthWhen growing materials, there are often two grow approaches: to grow us-ing a codeposition or to grow using a shutter method. Codeposition refersto when all elements are simultaneously deposited and shutter growth is adeposition method in which each element is deposited individually. In thecodeposition method, it may be hard to control the type of growth mech-anism. If there is no self assembly of the crystal structure, it is possibleto have a mess of elements with the correct stoichiometry form multiple273.1. Sapphire FGT growthphases. This is further complicated when involving materials with high va-por pressure, as it is not necessarily safe to assume the sticking coefficient ofelements or compounds (the relative number of atoms deposited as comparedto the number of atoms which strike the surface) while depositing at hightemperature, because of their temperature dependence [49]. In practice, itis not straightforward to determine the stoichiometry. Another difficultyarises when having to calibrate fluxes of each elements to agree within a fewseconds. Each flux calibration averages roughly 30 minutes. Thus, tuningthe deposition rates precisely can be a multiple hour task.On the other hand, in the shutter method, it may be hard for elements tofind their home and have the proper mobility, as each element is depositedat cooler temperatures as a way to control the SC. One also has to nav-igate unwanted phases growing before all elements have been introduced,especially when there are 3 unique elements all being epitaxially delivered.In the early trials, codeposition was the method of choice. Fe, Ge and Tefluxes were measured with a quartz crystal monitor (QCM) and stoichimet-rically calibrated. Fe, Ge and Te deposition rates were set such that a singlelayer or half unit cell of FGT was to be delivered to the substrates surfaceover a period of 580 seconds. The substrate temperature was held at 340◦C,the Fe, Ge and Te crucibles were held at 1120, 997, and 315 ◦C respectively.After 3 layers(1.5 unit cells) the substrate RHEED pattern was no longervisible and the new surface had significantly less RHEED intensity, almostappearing amorphous as displayed in Figure 3.8.283.1. Sapphire FGT growthFigure 3.8: RHEED image of the first attempt at growth of FGT usingthe codeposition method on the sapphire surface. Three streaks are faintlyvisible. In-plane lattice parameters were extracted to be roughly 4A˚. Note:the lack of general intensity is due to a low RHEED emission current andsignificant amount of diffuse scattering.293.1. Sapphire FGT growthFigure 3.9: Line profiles of a) sapphire substrates and b) possible films.Left: RHEED images of the grown film’s [112¯0] direction. Right: corre-sponding line profiles of the diffraction pattern. Lattice constants are ex-tracted using the ratio of the pixel peak positions. Note: the poor intensityin b1) is likely due to disorder and low emission current.A line profile was taken of the two RHEED directions of Figure 3.8 andused to determine the in-plane lattice constants of what was grown. Uponinvestigation, the in-plane lattice parameters agreed well with those of FGT.The theoretical mismatch was calculated to be 15.7% and the experimentalmismatch was found to be 15.8%. This film was grown only a few layers soit did not permit for further diffraction experiments like XRD, but yieldedpromising in-plane lattice parameters as detailed in Figure 3.9Future attempts at the growth of FGT were performed with the shuttermethod, in which subsequent layers of Te, Fe, FeGe, Fe and Te were de-posited in that order. Each layer was deposited at 200 ◦C. with the samesubstrate preparation as described above. Early attempts resulted in poly-crystalline films as shown in Figure 3.10 where the circular diffraction ringsindicate no preferred orientation of the crystal.To gain more information on the grown materials, lattice constants bothin- and out- of plane were analyzed using line profiles as displayed in Figure3.11. In this sample, a Fe:Ge:Te stoichiometry of 2.7:1:2 was used. TheFe ratio was chosen such that no excess Fe was deposited, but still withinthe window of stable FGT Fe concentration [24]. The Te flux was chosento match that of bulk crystal growths The film did not display a change in303.1. Sapphire FGT growthFigure 3.10: RHEED images of a polycrystalline film grown on the sapphire(0001) surface. Polycrystallinity is inferred from the diffraction rings.RHEED upon rotation of the sample indicating no preferred orientation ofthe film. The film also displayed polycrystalline rings. This made it impos-sible to determine the crystal symmetry. Yet, by analyzing line profiles ofthe RHEED images, it was possible to determine in-plane lattice constants.The lines are displayed in Figure 3.11 a1) and b1) by the thin lines. Thelattice constants were analyzed by the distances between peak as indicatedb y the black lines in all sub-figures. The in-plane constant was determinedto be 3.66 A˚and the c lattice constant was determined to be roughly 8.3A˚which are in disagreement with expected FGT lattice parameters.313.1. Sapphire FGT growthFigure 3.11: RHEED image of polycrystalline sample with 3D character.Line profiles of both the in-plane and out-of-plane lattice constants are dis-played in a and b respectively. The c lattice constant was determined tobe 8.3 A˚The in-plane lattice constant was determined to be 3.66 A˚. Thesesamples were grown with a 2 x over flux of Te atoms.Further attempts were made by changing the Te flux ratio. SC’s wereassumed to range anywhere between 0.75 to 0.5 as displayed in Figure 3.12.Each film resulted in a polycrystalline, three-dimensional diffraction pat-tern. The diffraction pattern of each sample did not change upon rotation.This indicates that the films were grown such that the film did not havea preferred orientation. Films with less tellurium tended to be more poly-crystalline while films with more tellurium tended to have a preferred orien-tation. Experiments to determine a more accurate telluirum SC were thencarried out. Samples were grown at 280◦C with up to 2 x Te overflux withcorresponding data displayed in Figure 3.12. This experiment highlightsthe necessary Te overlux required in the growth of FGT. As the Te flux isincreased the film began to become less polycrysatlline and 3D. Once thetellurium SC had been assumed to be 0.5 growths were more controllable.At this point the substrate temperature was varied to find the optimal depo-sition temperature with little success. Thicker films of 20 layers were grownand diffraction experiments were carried out and are recreated in Figure3.13323.2. Ge FGT growthFigure 3.12: Examples of RHEED images with varying amounts of Te fluxin codeposition growth. Substrate temperature was held at 300C. a) theSC is assumed to be 0.75 thus 3 of every 4 atoms were assumed to stickto the film surface. The surface appear very three dimensional and slightlypolycrystalline. b)SC’s assumed to be 0.6 resulting in a less polycrystallinesample but still 3 dimensional. c) 2 dimensional streaks are visible with asome 3 dimensional spots. All emission current was measured at 0.2µAIn codeposition growths at 300◦C the diffraction patterns began to changeas the samples were rotated indicating at least a preferred film orientation.The RHEED appeared streaky indicating smooth surfaces, and the in-planelattice constants were in agreement with FGT. Yet the films RHEED showedthat the growth was still not perfected. The existence of a double streakmarked with blue stars at what was assumed to be the 2nd order diffractionstreak indicated either multiple phases (the existence of what could be the(10) diffraction peak of FeTe,or a double domains [112¯0] direction overlapwith the [101¯0] direction). Diffraction experiments revealed there were mul-tiple phases of FeGe compounds, but also the existence of FGT. Furthertrials of the growth of FGT on sapphire were unsuccessful.3.2 Ge FGT growthCodeposition vs Shutter GrowthDue to the difficulties of growth on the Al2O3 surface, a different substratewas used. Ge (111) was chosen due to its matched in-plane lattice symmetryas well as its similar in-plane lattice constants of 4.013 A˚ resulting in a 0.5%lattice mismatch. This substrate also granted us the ability to deposit a Gewetting layer on the Ge surface to ensure atomic smoothness as well as bury333.2. Ge FGT growthFigure 3.13: First truly 2 dimensional RHEED images obtained duringshutter growth on sapphire.a) and b) were taken upon 30 degree rotationwhich would reveal the two unique crysallogrpahic directions. The RHEEDprofile did not change upon 30 degree rotations indicating a film with nopreferred orientation. Double streaks were also observed at the 2nd orderdiffraction points which indicated either multiple phases of Fe, Ge com-pounds or the incident beam probed areas with different FGT orientation.RHEED emission current was measured to be 0.3 µA upon each image taken.These films were determined to be significantly Te deficient.any contaminants that may exist on the substrate surface after placementinto the MBE chamber.Like many substrates, Ge is prone to surface contamination. In particu-lar, a native oxide is known to form on the surface of Ge(111). Furthermore,organic contaminants likely reside on the sample surface which are not eas-ily removed. Thus an etching procedure was developed to rid the substratesurface of organic compounds and remove the surface oxide layer.Delivered samples were subject to a “piranha etching” procedure to re-move the organic compounds. This involved a mix of 1:3 solution of hydrogenperoxide (H2O2) to sulfuric acid (H2SO4). It is important to note hydrogenperoxide was added after sulfuric acid for safety reasons. The solution wasthen heated to 130◦C on a hotplate while stirred. It was soon realized thatthis process was very time sensitive. Substrates left in the solution for toolong would develop a unidentified milky haze on the surface which was notremovable, thus rendering the substrates useless. Optimal time in the pi-ranha etch was determined to be roughly 30 seconds while on the hotplate.The substrates were subsequently placed in deionized water(DI), then addedto a second solution of 23:1 DI water to hydrofluoric acid (HF) for 30 sec-onds. Samples were again placed in DI water until loaded into the chamber.Substrates were finally degassed in the MBE chamber. Temperatures wereramped at 20 ◦C/min to 480 ◦C while maintaining chamber pressures be-low 5x10−9 torr. Upon approach to 480 ◦C the Ge(111) surface obtained a343.2. Ge FGT growth2x2 reconstruction. The reconstruction did not survive temperatures muchlarger than 480 but were robust while cooling. Once the reconstruction wasobtained, two additional monolayers of Ge were deposited or for as long asthe RHEED intensity increased.Examples of the substrate at 200◦C and at 480◦C are displayed in Figure3.14. These images were taken in the two unique crystallographic directionsnamely the [101¯0] and the [112¯0] directions labeled a) and b) respectively.At 200◦C the Ge surface does appear crystalline but the RHEED intensity islargely diffuse. This is likely due to residual surface contamination developedwhile transferring the substrate from the DI to the chamber. The intensityof the spectral and higher order diffraction spots observed in the highertemperature images a2) and b2) are indicative of a highly ordered atomicallysmooth surface. The less intense spots located between the bright spots arecharacteristic of a 2x2 reconstruction. After the reconstruction, substrateswere placed in a diffractometer to determine the exact position of Ge peaks.XRD was performed in a Rigaku Smartlab system at room temperature.A parallel beam (PB) configuration was implemented with a copper anodewhere both Kα1 and Kα2 emission channels were used. The voltage andcurrent were set to 40 kV and 44 mA, respectively.A Ge sample was first inserted into the instrument to note the peaksbelonging to the substrate. A 2θ-ω scan is displayed in Figure 3.2. The(111) and (333) peaks are most prominent at 27 and 90 degrees respectively.The (311) peak is also observed at roughly 57 degrees. the (111) peakcorresponds to a lattice spacing of 3.27 A˚.353.2. Ge FGT growthFigure 3.14: RHEED images of Ge samples in the two unique crystallo-graphic directions. a) and b) are Ge samples in the 101¯0 and [112¯0] di-rections respectively. a1) and b1) are taken at low temperature (200◦C)while a2) and b2) are taken at 480◦C. The streaks not visible in a1) andb1) are indicative of a reconstruction, more specifically since the half orderstreaks are observed in both crystallographic directions the reconstructionis determined to be 2x2.Like the sapphire substrates, both shutter and codeposition methodswere performed with varying degree of success. Shutter method growthswere attempted with substrates held at room temperature. The Fe, Geand Te cell temperatures (deposition times) were held at 1180(240s), 1070(130s) and 320 (117s) ◦C respectively. Because the QCM calibrations wereperformed at room temperature, the Te SC with respect to QCM was de-termined to be 1. This was a major factor in deciding to deposit at roomtemperature. The shutter growth method was performed layer by layeraccording the structure of the material thus stochimetric ratios of Te, Fe,FeGe, Fe, Te layers were deposited subsequently. Corresponding RHEEDimages are pictures in Figure 3.16, where a 10 unit cell (UC) film is pictured.The sample appears single phase in the RHEED images with in-plane latticeconstants corresponding to 4 A˚with no polycrystallinity.363.2. Ge FGT growthFigure 3.15: 2θ-ω scan from 0-100 deg. of the Ge(111) surface. Note thereare three main peaks at 27, 90 and 57. Those at 27 and 90 deg are the Ge(111) and (333) peaks respectively corresponding to a lattice spacing of 3.27A˚ The less intense peak at roughly 57 deg. belongs to the (311) diffractionplane.Figure 3.16: RHEED images of FGT shutter method of 10 UC of FGTand other phases. a) RHEED beam aligned with the [112¯0] direction andb) RHEED beam aligned with the [101¯0] direction. The RHEED emissioncurrent was held at 0.3µADiffraction measurements were then attempted to determine the c lat-tice parameters of the film and to identify any other phases that may haveexisted. The diffraction scan in shown in Figure 3.17.The scan was preformed from 0 to 90 degrees in air at room temperature.373.2. Ge FGT growthFigure 3.17: XRD 2θ-ω scan from 0-90 degrees of shutter growth methodof a 10 UC film on Ge (111). The sample has multiple Te deficient phasessuch as Fe3Ge2. Yet there are signs of FGT as indicated by the FGT labeledpeaks.The sample was found to contain multiple phases, namely FGT and FGT.This was a sign that the corresponding film was Te deficient. The mostpronounced peak is the Ge (111) peak. This is due to the thickness of thesubstrate. The (0 0 2n) peaks of Fe3GeTe2 are observed indicating progressat the growth of single phase FGT. There are also other phases observed inthe scan namely the Fe3Ge2 phase.AFM experiments were then carried out to analyze the topography ofthe sample. At this point the effects of exposure to oxygen were still not wellknown. Nevertheless, AFM scans were taken revealing a very rough surfaceas displayed in Figure 3.18. The height variation over the 2x2 µm scan wasas large as 17 nm. This was a sign that growth at higher temperature couldbe necessary. The increased growth temperature could in theory lead tobetter film quality because of the increased mobility of exptaxially deliveredatoms, but also alter the Te sticking coefficient.Once some FGT phase was confirmed, it was then possible to performtransport measurements to probe for magnetic order. The Rxy scan of a 10UC FGT film grown by shutter growth is recreated in Figure 3.19. At lowtemperatures (2, 50 K) hysteresis are observed implying ferromagnetic orderat low temperature. The transport data revealed that the FGT phase yield383.2. Ge FGT growthFigure 3.18: 2x2 µm AFM image of a 10 UC FGT surface grown onGe(111). Sample was grown using the codeposition method. The surfaceappears quite rough. Height variations over the scan are as large as 17 nm.Later experiments would further complicate analysis since the FGT surfaceis highly sensitive to ambient conditions. This sample was exposed to suchconditions for multiple hours.a signal despite the existence of other phases. The FGT phase responsiblefor the data must have existed across the entire substrate.Following the shutter method growths resulting in rough surfaces, thecodeposition method was attempted at higher substrate temperatures withparameters as follows. An Fe:Ge:Te: stoichiometry of 2:1:4 was used with amonolyaer growth period of 240 s, and Substrates were held at 320 ◦C. Afterdeposition of 4 layers substrates were annealed to 340 C for 1 hour. Thisprocess was repeated until 10 UC had been deposited just as in the shuttergrowth process. As displayed in Figure 3.20 the growth led to a surfacewhich was two dimensional as indicated by the streaks nature of the RHEEDintensity. In Figure 3.20 a) and b) correspond to the two unique diffractiondirections namely the [112¯0] and [101¯0] directions respectively. The emissioncurrent held at 0.2 µA for both pictures. In either direction no secondaryphases or polycrystallinity was observed. This was a promising result, asno single phase codeposition had been observed in sapphire growths. Thesample was then fully characterized with XRD, AFM and transport.In the diffraction experiment, the sample revealed multiple peaks cor-responding to materials other than FGT. The same Fe3GeTe2 peaks wereobserved as in the shutter growth sample. But the relative intensity of these393.2. Ge FGT growthFigure 3.19: Transport data of FGT shutter method of 10 UC of FGT andother phases. Data was taken in a Quantum Designs PPMS. Out-of-planemagnetic fields were swept from -3 to +3 T at temperatures ranging from2 to 200 K. The closure of the hysteresis loop implies a transition from theferromagnetic to paramagnetic state at 50 K.Figure 3.20: RHEED images of FGT co-deposition method of 2 UC of FGTand other phases. a) RHEED beam aligned with the [112¯0] direction andb) RHEED beam aligned with the [101¯0] direction. The RHEED emissioncurrent was held at 0.2 µApeaks relative to that of shutter growth was much less. But, like shuttergrowth, the codeposition recipe was still Te deficient.Following in-situ characterization the film was taken the AFM in order403.2. Ge FGT growthFigure 3.21: 2θ-ω scan of a 10 UC co-deposited FGT film on Ge(111). Thedata reveals a mixed phase film with numerous diffraction peaks. Thosebelonging to FGT and Fe3Ge2 are labeled.to examine its typography, particularly the smoothness of the surface. Theresulting 2x2 µm scan of the surface is displayed in Figure 3.22. The verticalrange recorded was roughly 2.5 nm over the entire scan range implying uni-formity over the substrate surface. Upon comparison to the shutter growthmethod, the surface was an order of magnitude smoother yet it was still notatomically smooth. The lack of atomic smoothness could be from a roughsubstrate surface, poor growth technique or oxidation of the surface uponexposure to air. The substrate and film surface imply atomically smoothsurfaces in RHEED, thus we were led to believe that the surface exposureto air was responsible.Transverse transport measurements on the sample were carried out from413.3. Capping DecappingFigure 3.22: 2x2 µm AFM image of the FGT surface grown on Ge(111).Sample was grown using the shutter method. Over the scanned area, heightvariation up to 2.4 nm was recorded2 to 250 K subject to magnetic fields ranging from -3 to +3 T. The corre-sponding data is displayed in Figure 3.23. The co-deposited sample showsan anomalous Hall response up to 250K. FM order indicated by a hysteresisloop halts at roughly 150K as indicated by the closure of the hysteresis.It was concluded from these experiments that the most promising ap-proach synthesizing atomically smooth single phase FGT was the codeposi-tion method.3.3 Capping DecappingAfter previous reports suggesting samples were air sensitive [50], a cappingprocedure was necessary to measure thin (few QL) samples. The cappingmaterial needed to be removable at low temperatures. More specifically,the cap needed to be removable at temperatures lower than the growthtemperature of the films in order to avoid sample degradation. Additional423.3. Capping DecappingFigure 3.23: Transport data of 10 UC FGT grown via codeposition withother phases. Out-of-plane magnetic fields were swept from -3 to +3 T attemperatures ranging from 2 to 250 K. The closure of the hysteresis loopimplies a transition from the FM to paramagnetic state at roughly 150 K.restrictions were imposed by the possibility of intermixing at the sample/capinterface. If the cap and the Te layer share the same valence states and aresimilar in ionic radii, it is likely that substitution would take place at theinterface. Because there existed no way to check for such substitution atUBC, it was in our best interest to stay away from materials which this waspossible.Ideally, the sample cap would also be amorphous to best protect thesample, as domain boundaries and other crystal defects invite oxidation andcontamination through such defects. Many of the elements available in GM2were out of the question. For example, Fe and Ge would almost certainly beirremovable due to their low volatility at low temperature. Compound capsgrown with multiple elements were a possibility, but all compounds availablesuch as GeTe and FeGe yielded lower vapor pressures than some individ-ual elements alone. At our disposal there existed two elemental candidates,namely Te and Se, which could serve to cap. Se has a higher vapor pressurethan Te and its volatility makes it easier to remove. Yet with Se depositionwe risk Se substitution at the surface/interface. Although this substitutionmay only be minor in monolayer and few layer experiments, these substi-tutions may have dramatic effects especially when the Te is responsible for433.3. Capping Decappingmany of the interesting physical phenomena that make FGT so attractivethrough means of spin orbit coupling.Te on the other hand also has a relatively high vapor pressure and anyintermixing of the cap and film was nothing to fret. Capping experimentswere carried out with Se and Te. First Te was attempted at room temper-ature as displayed in Figure 3.24 where a) is the diffraction pattern of ahighly ordered monolayer FGT sample which is indicated by the point-likediffraction. In b, a 20nm Te cap is displayed. During the capping proce-Figure 3.24: RHEED images of FGT film and cap. a) The RHEED profileof FGT film along the [112¯0] direction with point-like diffraction spots sig-naling a well ordered atomically smooth film.b) Te cap indicating 3D surfacedefects.dure the Te cell was held at 320◦C for 30 minutes. The film was held atroom temperature. The resulting cap was initially amorphous, but duringthe growth become crystalline and slightly 3D as indicated by the spottyand streaky nature of b). Decapping of the sample was then attempted.The sample was heated from room temperature at a 5 ◦C/min ramp rate.At 250 ◦C the sample was held for 30 minutes with no change in RHEED.After 30 minutes the sample was heated to 280 ◦C with a 1 ◦C/minute ramprate. At this point the cap started to decompose, and annealing was held at280 ◦C, the surface was left resembling that of FGT as displayed in Figure3.25. The resulting RHEED intensity along the [112¯0] direction was largelydiminished with a significant amount of diffuse scattering. The in-planelattice constants were similar to those of the original film. The samplewas then removed from the chamber where diffraction measurements wereperformed in the usual manner. The XRD scan revealed no FGT phasesleft therefore indicating that the decapping process was detrimental to theFGT film. This was interesting because the sample was never raised above443.3. Capping DecappingFigure 3.25: RHEED image of film post Te decapping experiments. Thediffraction pattern slightly resembles the [112¯0] direction of FGT. Emissioncurrent is held at 0.2 µA. Note: the blur seen in the image is due to anincrease amount of diffuse scattering due to a imperfect surface.the growth temperature during the decapping process.Figure 3.26: RHEED images of a) 2 UC FGT sample and b) amorphousSe cap.After the failed Te capping process a Se cap was attempted. The samethickness of cap (20nm) was deposited on on a 10 QL FGT sample as dis-played in Figure 3.26. The crystal quality of the film was very good, withpoint-like diffraction and a negligible amount of diffuse scattering. The Secap also grew quite well. The resulting cap was completely amorphous asindicated by the lack of any concentrated intensity in the image. The de-capping process was attempted in a similar fashion to the Te cap, but the453.3. Capping Decappingramp rate differed slightly. The Se decapping process was started at roomtemperature and was then heated at 2 ◦ per minute to 250 ◦C where theRHEED intensity began to change. Initially a diffraction pattern resemblingFGT was observed but with a 2x1 reconstruction as displayed in Figure 3.27.Figure 3.27: RHEED image of a film surface during the decapping proce-dure. the surface shows signs of a 2x1 reconstruction as indicated by thefaint half order streaks.After 10 minutes at 250 ◦C the sample began to become 3D, and wors-Figure 3.28: RHEED images of the Se decapped FGT films. The two uniquecrystallographic directions both exhibit 3D surfaces with what appears tobe a slight 2x2 reconstruction. The reconstruction was not removable underany conditions which conserved the underlying film. Thus proving Se wasnot an ideal candidate for capping/decapping procedures.ened quickly. The symmetry of the crystal remained 3 fold and the twocrystallographic directions are shown in Figure 3.28. Interestingly, the c463.3. Capping Decappingaxis parameter was not shared between the two directions which is illus-trated by the intensity modulations following a specific streak vertically.Due to the obvious degradation of the sample no further characterizationmeasurements were performed. From these experiments the hopes of suc-cessful decapping procedures were abandoned. This implied that any surfacesensitive measurements were required to be performed in-situ to avoid ex-posure to the atmosphere. Once a cap was deposited the cap would needto remain on the sample surface to ensure pristine films. It then impliesthat the techniques like scanning tunneling microscopy and angle resolvedphotoemission spectroscopy would be out of the realm of possibilities unlesssuch apparatuses were to be directly connected to the system or a vacuumsuitcase was used.47Chapter 4Experimental results4.1 Structural Characterization of Pristine FGTSamples by Co-DepositionIn early experiments co-deposition seemed to be the be the most promisingroute as indicated by transport and diffraction measurements, thus the sub-strate temperature and Te over flux parameter space was explored. Growthabove 360◦C was found to decompose FGT phases, and growth below 280◦Cled to FeTe phases as indicated by diffraction measurements. Sample growthtemperatures were chosen to be 330◦C at which Te overflux was optimizedat 320% where up to 400% Te flux did not change the growth, thus one couldassume that the constant deposition was in equilibrium with evaporation atthe sample surface. Higher than 400% excess Te was shown to decrease theRHEED intensity such that the evaporation rate could not match that ofthe deposition rate.Source materials of Fe, Ge, and Te were evaporated from the effusion cellsat 1320◦C, 1100◦C, and 325◦C, respectively. Elements were co-deposited onthe substrate at 330◦C to grow FGT. The ratio of Fe:Ge:Te was set at 3:1:5.6 measured by QCM prior to growth. These temperatures and flux ra-tios were determined to yield optimal FGT films. It was then possible togrow optimized high-quality growth of epitaxial FGT(0001) ultra-thin filmswith thicknesses ranging from 1 QL to 10 QL on atomically smooth Ge(111)substrates. Corresponding FGT RHEED images are displayed in Figure4.1. a) and b) are the usual Ge surfaces where the 2x2 reconstruction isapparent as displayed by the half order streaks. c) and d) are images of a4 QL FGT surface where point-like diffraction spots are indicative of wellordered atomically smooth surfaces. Both crystallographic directions aredisplayed in Figure 4.1 in which a hexagonal lattice directions , to whichFGT and Ge(111) belong, there exist two distinct crystallographic direc-tions separated by 30 degrees, the [112¯0] and [101¯0] directions. In-planelattice parameters are measured in-situ throughout the growth process us-ing RHEED line intensity profiles of the associated diffraction patterns. The484.1. Structural Characterization of Pristine FGT Samples by Co-Depositionlattice constant of the FGT film 3.99 A˚agrees with the known in-plane lat-tice constant for bulk Ge(111), as well as the bulk FGT measurements [32]It is essential to note the diffraction patterns of the two crystallographicdirections of the substrate and film are aligned, which confirms an epitaxialgrowth as displayed in Figure 4.1.Figure 4.1: In-situ RHEED images of a 22 reconstructed Ge (111) sub-strates along the [101¯0] and [112¯0] direction. Bottom: 4-QL Fe3GeTe2 withelectron beam along [101¯0] direction and [112¯0] direction.A layer-by-layer growth mode is established via observation of RHEEDintensity oscillations as displayed in Figure 4.2. The oscillation period allowsthe precise control of the number of layers deposited in thick samples as wellas determines the exact amount of time-per-layer to minimize layer number494.1. Structural Characterization of Pristine FGT Samples by Co-Depositiondifferences within a sample. The oscillation period was determined to be78 seconds which was highly reproducible. In thinner samples a calibrationsample was used to determine the layer oscillation period such that shutterscould be closed before the downturn in the intensity and correspondingly thenext layer began to form was observed. In this experiment the 1st diffracted(01) and specular spot (00) in Figure 4.1 oscillations were in-phase, and theoscillations began at a maximum thus we could conclude that a downturnin the intensity indicated the start of a new layer.Figure 4.2: RHEED intensity oscillations during the growth of Fe3GeTe2,indicating atomic smoothness and its epitaxial growth mode. Single QLgrowth times are extracted from the period of oscillation to be 80 s.The lattice structure of FGT films were further characterized by ex-situ XRD. A single phase of FGT (001) films is confirmed by the (002)family of diffraction peaks in Figure 4.3. The diffraction pattern suggestsa out-of-plane lattice constant c=16.355 A˚, in good agreement with bulkmeasurements of 16.33A˚[32]. Fe vacancies in the Fe2+ site have been shownto decrease in-plane lattice constants and increase c [51], which are not ob-served in our films. Experiments were performed to conclude if the FGTsurface is air sensitive as no previous reports of FGT did detail cappingprocedures[24], while also stating fabrication had been done in oxygen freeenvironments. A 10 QL FGT sample was removed from the MBE cham-504.1. Structural Characterization of Pristine FGT Samples by Co-DepositionFigure 4.3: Ex-situ XRD of the out-of-plane lattice structure of FGTfilms. A single phase of FGT (001) films is confirmed by the (002) familyof diffraction peaks. The diffraction pattern suggests a out-of-plane latticeconstant c = 16.355 A˚, in good agreement with bulk measurements of 16.33A˚. The shape of peaks has been observed for all samples of pure FGT, theodd shape of such peaks could be due to thin samples unable to generatesignificant peaks or large scan speeds.ber and exposed to ambient conditions for a few hours. Once placed backinto the chamber the surface had become completely amorphous indicatingno crystalline structure (amorphous) remained after exposure to ambientconditions. Upon placement in the chamber, the sample was slowly heated(softly baked) to 340◦C in at a rate of 2 deg/min while chamber pressureswere kept below 5x10−9 torr. The surface did not recover in any soft bakeexperiments. An example of such a surface is displayed in Figure 4.4514.2. Transport results showing the conductivity and ferromagnetism in Fe3GeTe2 filmsFigure 4.4: Diffuse scattering FGT sample after exposure to ambient con-ditions for only a few minutes. The lack of features reveled by the RHEEDimage implies an oxidized amorphous surface. This is hard evidence thatthe surface layers of FGT are air sensitive. The amorphous layers were notremovable after soft bakes up to 340C. This also served as an indication thatany AFM scans are not of the FGT surface but of an oxidized surface thatmay not reflect the surface smoothness of the film in the chamber.4.2 Transport results showing the conductivityand ferromagnetism in Fe3GeTe2 filmsAfter fully characterizing the structure, single phase single domain FGTfilms could be measured in a PPMS, where electronic and magnetic prop-erties of FGT films were characterized by transport measurements. Waferscale measurements were possible with no lithography involved as pictured inFigure 4.5. Gold wires were indium cold pressed directly onto the sample(Tecap) surface, and to the sample puck pads. Samples were subsequently in-serted into the PPMS.Longitudinal and transverse resistances Rxx and Rxy (Figure 4.6) were524.2. Transport results showing the conductivity and ferromagnetism in Fe3GeTe2 filmsFigure 4.5: A photo of an FGT sample measuring the longitudinal resis-tance with van der Pauw geometry in a PPMS .measured on both capped and uncapped samples ranging from 1 QL to10-QL films. Here it is important to not that the remainder of data cor-responds to capped samples unless otherwise stated. Uncapped monolayerand bilayer films did not permit transport measurements due to insulatingbehavior, likely due to the oxidization of the FGT surface [50]. Differencesin transport behavior are negligible between capped and uncapped sampleswhen the thickness ≥ 5QL.Rxx measurements show resistance values con-sistent with exfoliated samples [24] and indicate a weak insulating behaviorat low temperatures in Figure 4.6. In control samples of Te capped Ge sub-strates (as shown by the purple line in Figure 4.6, we observe a strongerinsulating behavior and a much higher resistance than FGT films at lowtemperatures. Above 200 K, the conduction through Ge substrates is re-sponsible for the higher temperature transition in Rxx. The next transitionis due to the emergence of localized states with decreasing temperature.Below 200 K, Rxx fit the variable range hopping model in 2D as shownin Figure 4.7, where Rxx ∝ eT−1/3. This is in agreement with previoustransport studies [47], which could be attributed to the emergence of lo-calized electrons with decreasing temperature. These localized states likelyoriginate from the Fe 3d orbitals, which are also responsible for the itiner-ant electrons. The transition temperature (the value at which the hysterisisloop closes agrees well with previous reports of a heavy-fermion state (“lo-534.3. Temperature dependence of the anomalous Hall behaviorFigure 4.6: Longitudinal resistivity Rxx vs. temperature T taken of variousQL numbers. A Te capped substrate is also measured as a test sample tosee the intrinsic film and cap longitudinal behaviour.calized”) in bulk [36] and correlated electrons.4.3 Temperature dependence of the anomalousHall behaviorTo study the phase transition from the FM to paramagnetic state temperature-dependent RAH measurements were carried out on all samples from 5 K to225 K, as presented in Figure 4.9. Magnetic order was first characterized at5 K by transverse resistance (Rxy) measurements that host the anomalousHall (AH)All samples display clear hysteresis indicating ferromagnetic order. 5-QLcapped and uncapped samples are shown in Figure 4.8 d, revealing an al-most identical remanent (RR) and saturation resistance (RS). This impliesthat, in uncapped samples above 5 QL, the top 1–2 FGT oxide layers donot significantly affect transport properties of the underneath pristine FGT544.3. Temperature dependence of the anomalous Hall behaviorFigure 4.7: Fit of the Rxx with the Mott variable range hopping model,indicating the existence of localized states below 125 K.Figure 4.8: RAH vs. applied field of select QL numbers at 5 K showing theAH and hysteresis loops. The applied magnetic field H is perpendicular tothe sample surface and parallel to the c axis of FGT. All samples are cappedexcept the two samples marked as non-capped (NC).554.3. Temperature dependence of the anomalous Hall behaviorlayers, suggesting the electronic channels are not confined to the surface.This result is significant considering previous reports of antiferromagneticcoupling between oxidized and pristine FGT, where one would expect un-capped samples to display a two-part hysteresis. We do not observe such asuperposition in hysteresis shape. A hole carrier type was determined fromRxy at all temperatures. Strong perpendicular magnetic anisotropy (PMA)indicated by the observed textbook-like square shape of the hysteresis. It isimportant to note that PMA and hysteresis width are proportional [52], andis also indicative of a single domain [52], where all Ising like spins coherentlyflip at the coercive field (HC) along the easy axis.Figure 4.9: Anomalous Hall resistance of a Single-QL, b 2-QL, c 5-QL,and d 10-QL FGT films. The hysteresis loop is apparent in all samplesat low temperatures and disappears as the temperature reaches the Curietemperature Tc564.3. Temperature dependence of the anomalous Hall behaviorFrom these evolutions we extract various quantities including the coer-cive field and the remanent and saturation resistance of samples at varioustemperatures (Figure 4.10). The coercive field (HC) is defined as the valueof the applied magnetic field at which the loop’s curve crosses through RAH= 0. The remanent and saturation resistance RR and RS are defined as thevalue of RAH at zero applied field and upon saturation, respectively. It iswell established that RR and RS are directly proportional to the magneticremanence MR (the magnetic signal at zero applied field) and magnetic sat-uration MS (the magnetic signal when all magnetic moments are maximallyaligned) [52]. To investigate the layer and temperature-dependent evolutionof the magnetic phases, we compare RR, RS , and HC as a function of layernumber at a fixed temperature. With increasing layer number, all sam-ples decrease in the remanent MR and saturation magnetization MS Figure4.10 at a fixed temperature below the respective Tc. This progression isconsistent with previous reports [24].The ratio of RR and RS is plotted in Figure 4.10 c. The ratio remainsnearly 1 for all QL numbers up to 100 K, where the ratios for 5- and 10-QLsamples exceed 0.9 at room temperature. We also note that the energy lossdecreases as the layer number is increased, as indicated by the area enclosedby the hysteresis loop after a complete magnetic field sweep.The field needed to reverse the direction of magnetization, known as thecoercive field Hc, provides valuable information for assessing the stability ofmagnetic structures and is often used as an indication of the hardness of themagnetic domains Figure 4.10 d). HC in monolayer samples is greater than1 T and roughly 0.5, 0.8, and 0.6 T for 2, 5, and 10 QL, respectively. Bilayersamples do not fit into a monotonic decrease of HC with respect to layernumber. One possible explanation is related to hole doping by Fe vacancies,which has been shown to decrease the perpendicular anisotropy [53]. A morethorough layer dependent study is required to reveal the underlying cause.The spin flip transition shown by the gradual progression of RAH around thecoercive field could be due to a lack of magnetic anisotropy or the existenceof multiple magnetic domains.Tc is determined by two methods. The remanent resistance RR serves asa direct probe of spontaneous magnetic order, MR, and allows us to extractTc of samples by tracking the temperature at which RR vanishes (Figure4.10 f). Tc monotonically increases with layer number. For 1, 2, 5, and10-QL transitions take place at roughly 75, 125, 150, and 175 K, respec-tively as compared to 220 in bulk [33]. A secondary method to extract Tc isby Arrott plot analysis [54] (Figure 4.10 e). Arrott plots employ the equa-574.3. Temperature dependence of the anomalous Hall behavior!hFigure 4.10: a) Remanent resistance (RR) and b) saturation resistance (RS)plotted against QL number at various temperatures. Both exhibit a mono-tonic decrease with increasing layer numbers. c) RR/RS ratios remain nearunity at different temperatures, suggesting FGTs viability in applications.d) Coercive field (Hc) of various QL layer numbers. e) Arrott plot analysisof a 5-QL sample, indicating a Tc of 176 K extracted by the x-interceptof the fitted line. f) Magnetic phase diagram of FGT as a function of QLnumber. Tc on the RAH method line is chosen at the lowest temperatureat which RR vanishes in AH measurements. Arrott plot points are foundusing methods as in e). The green region between the paramagnetic andferromagnetic state are unmeasured regions where the ferromagnetic statecould still survive.584.3. Temperature dependence of the anomalous Hall behaviortion RAH = µ0H/RAH , from which the high field portion is linearly flitted(inset of Figure 4.10 e). The intercept of RAH is then found for multipletemperatures and again linearly fitted as a function of temperature. Theresulting temperature intercept is then used as Tc. This method yields anexact temperature for Tc as compared to the range generated from RAHmeasurements. Using this method, we have extracted Tc to be 126 K inmonolayer samples, 145 K in bilayer samples 176 K, and 178 K in 5 and10-QL samples, respectively. The two methods show different Tc in samplesless than 10-QL, but are in agreement with RAH measurements of 10-QLsamples. Upon comparison with the exfoliated layer dependant phase dia-gram of Deng et al., we see a significant increase in Tc for monolayer andbilayer samples (roughly 100 K increase), but good agreement for 5 and10-QL samples.59Chapter 5Discussion & ConclusionWe note that for all layer numbers, FGT retains a strong anisotropy as in-dicated by the shape of the hysteresis, implying that there is limited holedoping, which has been shown to decrease the magnetocrystalline anisotropyenergy and lead to small coercive fields [53]. This is corroborated by bulksample measurements, which also show strong perpendicular anisotropy [37].Ultrathin UHV-grown monolayer samples seem to be of significantly higherTc and more robust ferromagnetism compared to exfoliated samples, as indi-cated by Hc, RR/RS ratio, and the field at which samples saturate indicatedby a flat AH signal. We speculate this result is due to damages undergoneby the sample during the fabrication process in exfoliation techniques. Thisis most obvious upon comparison to monolayer hysteresis reported by [24],for example, the coercivity of our epitaxial samples is nearly an order ofmagnitude larger than exfoliated samples. For monolayer, magnetic satura-tion was not achieved past 10 T while saturation was reached at only 2 Tin exfoliated samples, yet bilayer coercivity samples agree. At low temper-atures, the magnetic remanence to magnetic saturation ratio is also muchlarger in epitaxial samples for layer number less than 5. Tc of monolayerand bilayer samples is higher than any previously reported using anoma-lous Hall measurements. Effects of domains could play a role in this result.Where the Tc could decrease due to the existence of domains developed tominimize the stray field. Unfortunately in our experiment domains werenot directly probed with either Lorentz transmission electron microscopy ormagnetic force microscopy thus we cannot make a definite remarks, althoughthe square shape of our hysteresis loop implies single domain across the en-tire sample. Considering the decrease in coercivity caused by oxidation,indicated by differences in 5-QL capped and uncapped samples, previousmonolayer data could be explained by a partially oxidized sample.The lack of apparent antiferromagnetic interlayer coupling should be dis-cussed with respect to recent reports of systems such as CrI3, and CrBr3. Itis well known that interlayer coupling is extremely sensitive to interlayer dis-tances, which allow for orbital overlap. The considerable distance betweenadjacent QL containing Fe of roughly 8 A˚ does not lend itself to strong60Chapter 5. Discussion & Conclusioninterlayer interaction, compared to systems such as CrI3 where the inter-layer distance is smaller (4.2 A˚). A recent study using immense pressuresrevealed that the antiferromagnet-ferromagnet switching is not a result ofthe interlayer coupling, but due to stacking order differences in CrBr3 [55].For FGT, recent DFT, neutron, and spin polarized scanning tunnelingmicroscopy studies [56] the interlayer exchange had been determined to beFM. They show that in calculations it is essential to include the hubbardU to correctly model interlayer exchange in FGT. In calculations whereU is not taken into account the AFM interlayer exchange is energeticallyfavorable, but in systems with nonzero U the FM phase is determined to bethe ground state. Furthermore, powder neutron experiments do not showdiffraction peaks associated with the AFM phase over a 300 degree windowindicating no change in magnetic order across the magnetic phase transition(FM to PM). Lastly, a spin polarized STM does not indicate changes in spindirection across a single layer step edge which would be expected in AFMinterlayer exchange.Looking forward, a recent publication shows the coupling of ferromag-netism and topological states introduces controllable currents in helical edgestates in CrI3 and WTe2 heterostructures [57], suggesting an interesting fu-ture in the heterostructure design of 2D magnetic materials. Future het-erostructure design could also be fabricated in techniques enabling controlof the relative stacking angle which may pose the possibility to control inter-layer interactions and larger superstructure Moire patterns leading to newand exciting spin textures [58].Planned experiments for FGT include determining the monolayer bandstructure probed via ARPES to determine the evolution form bulk to mono-layer and the search highly correlated electron behaviour in the existance offlat bands which have been theorized and seen in thicker samples in ARPESexperiments [36].In conclusion, we have synthesized high-quality ultrathin epitaxial FGT(001) films ranging from monolayer to ten layers. The ability to grow atthe wafer-scale with precise control of layer number makes production-scaledevice fabrication a possibility. 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Molecular beam epitaxy growth and properties of Fe3GeTe2 thin films
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