Integration of in situ RHEED with magnetron sputter deposition for

atomic layer controlled growth

By

Jacob P. Podkaminer

A dissertation submitted in partial fulfillment of

the requirements for the degree of

Doctor of Philosophy

(Materials Science)

at the

UNIVERSITY OF WISCONSIN‐MADISON

2016

Date of final oral examination: 04/19/2016

The dissertation is approved by the following members of the Final Oral Committee:

Chang-Beom Eom, Professor, Materials Science and Engineering

Mark S. Rzchowski, Professor, Physics

Robert F. McDermott, Professor, Physics

Xudong Wang, Professor, Materials Science and Engineering

Thomas Tybell, Professor, Electronics and Telecommunications

i

ii

Abstract

Integration of in situ RHEED with magnetron sputter deposition for

atomic layer controlled growth

Jacob P. Podkaminer

Work completed under the supervision of Professor Chang-Beom Eom

University of Wisconsin – Madison

Epitaxial thin films continue to be one of the most promising topics within electronic

materials research. Sputter deposition is one process by which these films can be formed and is a

widely used growth technique for a large range of technologically important material systems.

Epitaxial films of carbides, nitrides, metals, oxides and more can all be formed during the sputter

process which offers the ability to deposit smooth and uniform films from the research level up

to an industrial scale. This tunable kinematic deposition process excels in easily adapting for a

large range of environments and growth procedures. Despite the vast advantages associated with

sputter deposition, there is a significant lack of in situ analysis options during sputtering. In

particular, the area of real time atomic layer control is severely deficient.

Atomic layer controlled growth of epitaxial thin films and artificially layered

superlattices is critical for both understanding their emergent phenomena and engineering novel

material systems and devices. Reflection high-energy electron diffraction (RHEED) is one of the

most common in situ analysis techniques during thin film deposition that is rarely used during

iii

sputtering due to the strong permanent magnets in magnetron sputter sources and their effect on

the RHEED electron beam. In this work we have solved this problem and designed a novel way

to deter the effect of the magnets for a wide range of growth geometries and demonstrate the

ability for the first time to have layer by layer control during sputter deposition by in situ

RHEED. A novel growth chamber that can seamlessly change between pulsed laser deposition

and sputtering with RHEED for the growth of complex heterostructures has been designed and

implemented. Epitaxial thin films of LaAlO3, La1-xSrxMnO3, and SrRuO3 have all been

deposited by sputtering and shown to exhibit clear and extended RHEED oscillations. To solve

the magnet issue, a finite element model has been constructed to predict and avoid the deflection

of the electron beam in many geometries. Together, this creates the possibility for RHEED to

become a widely used real time analysis tool with sputter deposition with far reaching

applications and potential.

iv

Acknowledgements

I will never forget the day that I was driving through the Rocky Mountains with friends

heading for Vail when, I received a phone call from Professor Chang-Beom Eom. He was

calling to make me an offer for a Research Assistant position in his lab and I had zero bars of

service. I knew this could easily be one of the most important calls of my life and naturally I

was panicking. In the end, I moved to Madison to work for Professor Eom and he has provided

me with continuous and steadfast support since. I cannot thank him enough for the number of

doors he has opened for me and the opportunities he has given me. In many ways he has enabled

the child inside me with the ability to play with all the toys in the lab and gave me the freedom to

get involved in so many projects when my mind wanders. It has been a truly great experience

and built a relationship that will continue for many years.

My fellow co-workers and collaborators, past and present, deserve a standing ovation as

well. Throughout the years they have helped me, pushed me, and put up with me. So many

great friendships and bonds have been built making this experience truly outstanding. In

particular, I would like to thank Chad Folkman for being my mentor early on and still, Camilo

for being a great housemate and research partner, Wittawat and Josh for going through all the

years of PhD work together and supporting each other, and Sangwoo who was my officemate,

mentor, friend, and collaborator. Bruce Davidson with whom I spent many a nights at several

synchrotrons, contributed to much of this work and is now a good friend. The entire Oxide Lab

deserves recognition and I thank you all. Additionally, all my friends and experiences I have had

outside of the lab have made this time balanced and enjoyable. A recognition is due to all those

that have contributed to my “normal person life”.

v

My family has been my rock throughout all of this. The PhD process is full of peaks and

extremely deep valleys at times. My parents, Joel and Tina, my sister Annie, and all my

extended family, they have been there through it all! They seem interested when I describe my

research, and were understanding when I vented my issues. No matter the situation they were

there for me and have given me the tools to succeed in life from day one. So much of my

success I owe to them.

And finally, my soon to be wife, Gina Furlano deserves the biggest thank you of all. You

came into my life when I was at my lowest point and almost immediately carried me to the

highest peak. You have been there for me day in and day out with continuous words of

encouragement. Through the good times and the bad, happiness has always been found in your

company. You never cease to amaze me and inspire me with your strength, independence,

positivity, and balance. I cannot wait for what our future holds. Thank you.

-Dedicated to Jack Stanley Podkaminer & Chis Bond

vi

Table of Contents

Abstract ....................................................................................................................................... ii

Acknowledgements ................................................................................................................. iv

Table of Contents .................................................................................................................... vi

List of Figures .......................................................................................................................... ix

Preface ........................................................................................................................................ xi

1. Introduction ...................................................................................................................... 1

1.1. Motivation ............................................................................................................... 1

1.2. Oxide Thin Films .................................................................................................... 4

1.3. Deposition techniques ............................................................................................. 6

1.3.1. Molecular beam epitaxy (MBE) .............................................................. 8

1.3.2. Pulsed laser deposition (PLD) ............................................................... 10

1.3.3. Sputter deposition .................................................................................. 13

1.4. Reflection high energy electron diffraction (RHEED) ......................................... 16

1.5. Outline of thesis .................................................................................................... 21

1.6. References for chapter 1 ....................................................................................... 24

2. PLD-Sputtering Hybrid with RHEED .................................................................. 31

2.1. Introduction ........................................................................................................... 31

vii

2.2. System design ....................................................................................................... 35

2.2.1. Deposition chamber ............................................................................... 35

2.2.2. LabView Control Program ..................................................................... 42

2.3. Metal – Oxide Heterostructures ............................................................................ 44

2.4. Conclusions ........................................................................................................... 57

2.5. References for chapter 2 ....................................................................................... 58

3. Two-dimensional electron gas (2DEG) at the LaAlO3/SrTiO3 interface by

sputtering........................................................................................................................ 61

3.1. Introduction ........................................................................................................... 61

3.2. Film growth and structural characterization ......................................................... 66

3.3. Electric characterization ........................................................................................ 70

3.4. Conclusions ........................................................................................................... 73

3.5. References ............................................................................................................. 74

4. In situ RHEED during oxide sputtering ............................................................... 78

4.1. Introduction ........................................................................................................... 78

4.2. RHEED with magnetron sputtering ...................................................................... 80

4.2.1. Scattering due to gas .............................................................................. 81

4.2.2. Deflection due to magnetic field ............................................................ 85

4.3. RHEED intensity oscillations ............................................................................... 89

4.3.1. LaAlO3/SrTiO3 ....................................................................................... 91

viii

4.3.2. La0.7Sr0.3MnO3/SrTiO3 ........................................................................... 95

4.3.3. SrRuO3/SrTiO3 ....................................................................................... 99

4.4. Conclusions ......................................................................................................... 104

4.5. Reference for chapter 4 ....................................................................................... 105

5. Finite element modeling ............................................................................................ 110

5.1. Introduction ......................................................................................................... 110

5.2. Modeling parameters .......................................................................................... 112

5.3. Single and two gun 90° off-axis geometries ....................................................... 112

5.4. Two antisymmetric configurations ..................................................................... 120

5.5. Generic solution .................................................................................................. 122

5.6. Antisymmetric SRO growth ............................................................................... 125

5.7. Conclusions ......................................................................................................... 129

5.8. References for chapter 5 ..................................................................................... 130

6. Summary and future thoughts ............................................................................... 131

ix

List of Figures

Figure 1.1. The classic perovskite unit cell..................................................................................... 5

Figure 1.2. Common lattice parameters and strain states.. ............................................................. 7

Figure 1.3. A schematic of an oxide-MBE chamber ...................................................................... 9

Figure 1.4. Example of sputter processes ..................................................................................... 13

Figure 1.5. Schematic representation of RHEED intensity oscillations ....................................... 20

Figure 2.1. MgO surface after exposure to air .............................................................................. 33

Figure 2.2. Schematic layout of the "hybrid" deposition chamber ............................................... 36

Figure 2.3. The mirrored growth geometries ................................................................................ 37

Figure 2.4. Growth sequence static RHEED images .................................................................... 46

Figure 2.5. Structural characterization .......................................................................................... 50

Figure 2.6. Two step Al2O3 RHEED progression ......................................................................... 52

Figure 2.7. Surface images of Al2O3 and Re ................................................................................ 53

Figure 2.8. TEM analysis of the heterostructure ........................................................................... 54

Figure 2.9. XRD of the Al/Al2O3/Re trilayer ................................................................................ 55

Figure 2.10. Quality factor measurements of the trilayer ............................................................. 56

Figure 3.1. Layering structure in the LAO/STO heterointerface .................................................. 63

Figure 3.2. Carrier concentration as a function of number of unit cells ....................................... 64

Figure 3.3. Structural characterization of the LAO/STO heterostructure ..................................... 68

Figure 3.4. Surface topography of the STO substrate and LAO film ........................................... 69

Figure 3.5. Transport properties of the sputtered LAO/STO heterostructure. .............................. 70

Figure 3.6. Room-temperature conductive-AFM (c-AFM) switching ......................................... 72

x

Figure 4.1. The differential scattering cross section for Ar and O atoms ..................................... 83

Figure 4.2. Scattering comparison for Ar and O2 gas ................................................................... 84

Figure 4.3. Schematic of the growth chamber .............................................................................. 88

Figure 4.4. Sputter deposition effects on RHEED. ....................................................................... 89

Figure 4.5. RHEED oscillations during a LaAlO3 growth on a SrTiO3 substrate ........................ 92

Figure 4.6. Ex situ structural and surface analysis of the LAO/STO film .................................... 94

Figure 4.7. La0.7Sr0.3MnO3 growth on a SrTiO3 substrate showing clear RHEED intensity

oscillations ........................................................................................................................ 96

Figure 4.8. X-ray and AFM measurements of the LSMO film .................................................... 98

Figure 4.9. RHEED oscillations from SRO growth .................................................................... 100

Figure 4.10. Topographic and structural characterization of the SRO film ................................ 102

Figure 5.1. 2-dimensional cross-sections showing the magnetic field close to the sample for the

single gun, two gun symmetric and two gun antisymmetric magnet polarities .............. 114

Figure 5.2. Finite element simulations showing electron beam deflection in single gun,

symmetric, and antisymmetric sputter source geometries and their resulting RHEED

pattern ............................................................................................................................. 116

Figure 5.3. Finite element simulations side view ....................................................................... 118

Figure 5.4. Magnetic field simulations for the two possible antisymmetric configurations ....... 121

Figure 5.5. Confocal and high rate off-axis antisymmetric setups ............................................. 123

Figure 5.6. An on-axis sputter arrangement with symmetric sputter sources centered around the

substrate .......................................................................................................................... 124

Figure 5.7. RHEED intensity oscillations during SRO growth .................................................. 127

xi

Preface

The work presented in this thesis has been the culmination of the author’s research over

the last six years. This work is primarily orchestrated and conducted by the author but naturally

done in collaboration with several other groups. The author has performed all of the film

growth, structural and surface characterization, and worked closely with J. Patzner to create the

finite element models. Chapter 2 is both original and work that has been adapted from APL

Mater. 1, 042115 (2013). The TEM in Chapter 2 was performed by Y. Zhang and X.Q. Pan.

The SSQCD devices and microwave characterization in Chapter 2 were performed by U. Patel,

Y. Gao, and R. McDermott. The films for the SSQCD were grown together with K.H. Cho. The

Al/MgO/Re trilayer XRD measurement was done by C.M. Folkman who also helped build the

hybrid chamber. The electrical characterization in Chapter 3 was done in collaboration with T.

Hernandez and M. Rzchowski, and the c-AFM work was done by M. Huang and J. Levy.

1

1. Introduction

1.1. Motivation

The focus of this thesis is on pushing the boundaries of thin film deposition in order to

create new material systems and their potentially novel electronic states. As Richard Feynman

said in 1959:

But it is interesting that it would be, in principle, possible (I think) for a physicist to

synthesize any chemical substance that the chemist writes down. Give the orders and the

physicist synthesizes it. How? Put the atoms down where the chemist says, and so you

make the substance. The problems of chemistry and biology can be greatly helped if our

ability to see what we are doing, and to do things on an atomic level, is ultimately

developed – a development which I think cannot be avoided.1

This prophecy has been a strong motivation for this thesis. While there are many interesting

material systems present in nature, the ability to create “artificial” materials can open the door

for new and exciting phenomenon to be discovered. While Feynman’s prediction has become

reality to a certain degree, there is always further development, which is what we aim to do.

Oxide materials, in particular the perovskite family, are an excellent place to pursue the

enhancement of atomic level control and the creation of new materials due to the nature of the

2

ABO3 crystal structure. While there are some restrictions, in general the A and B constituent

elements are freely interchanged to access new materials properties. This gets even more

exciting when combing two distinct perovskites together, which in general is feasible due to the

same parent crystal structure and oxide nature. The exciting physics and material science that

can be discovered and engineered at the materials interface and the coupling between their

discrete electronic properties is what drives this field in general.

To obtain this type of scientific exploration, perfect intrinsic materials are desirable, free

of defects and extrinsic contaminates. Epitaxial thin films provide the ideal platform for

developing new and exciting electronic materials for several reasons. First, thin film deposition

brings us closer to Feynman’s idea for creating new materials since it is a growth process that is

on the nanometer to sub-nanometer scale already. Second, most epitaxial deposition processes

take place in controlled vacuum environments, which limits the exposure to external

contaminants. In this way, interfaces can be kept as close to pristine as possible. Third, coherent

film growth facilitates the engineering of physical states into the material system which are

otherwise extremely challenging to obtain in bulk materials. An excellent example of this is thin

film epitaxial strain which can reach equivalent pressures in the range of GPa and can

significantly affect the material property such as increasing the superconducting Tc.2,3

And

finally, it is possible to deposit complex heterostructures and superlattices to build material

stacks that otherwise do not exist. This is not an exhaustive list of the advantages to epitaxial

thin film growth and while there are disadvantages and limitations compared to polycrystalline

or bulk materials, single crystalline thin films are a direct route to achieving and probing the

intrinsic properties of a material.

3

While epitaxial thin films are the perfect platform for pushing the bounds of material

properties, a way to “see” what is happening during the deposition is ultimately needed to garner

true atomic level control. For this type of in situ control, reflection high energy electron

diffraction (RHEED) is often utilized. Real time monitoring of the RHEED pattern provides

information regarding the interface state, lattice relaxation, and growth dynamics, and allows for

precise thickness control. In molecular beam epitaxy (MBE), RHEED is also commonly used

for sub unit cell thin film growth which is approaching the level that Feynman discusses.

Despite the popularity of RHEED, there is still room to improve. For instance, only recently has

pulsed laser deposition (PLD) shown that it can approach the type of atomic control that is

observed via RHEED in MBE.4 On the other hand, despite sputter deposition being a prominent

growth technique for scientific research and at an industrial level for a wide range of materials, it

rarely is used in conjunction with in situ RHEED.5 This has severely inhibited sputter

deposition’s ability to excel in the area of precise epitaxial oxide heterostructure growth to the

level that MBE and PLD has.

The lack of in situ analysis during sputter deposition is the primary motivation for this

thesis. The goal has been to determine precisely what limits the integration of RHEED with

sputtering, find ways around the limitations, and ultimately solve the problems to either remove

or mitigate issues to the point that they are not significant factors. Meanwhile, the goal has also

been to use RHEED with sputtering to grow interesting and novel heterostructures, from the

LaAlO3/SrTiO3 two-dimensional electron gas (2DEG) to epitaxial capacitive stacks for solid

state quantum computing devices (SSQCDs). Following the wisdom of Feynman, it is also of

interest for this thesis that this technique be used to further approach the ability to build any

material system desired by placing atoms where we want them. As such, the future of this area

4

of research is geared towards depositing materials that are outside the realm of possibility for

both MBE and PLD and require the precise control provided by in situ RHEED monitoring

during sputtering.

1.2. Oxide Thin Films

Complex oxides have some of the most diverse physical properties known to a single

class of materials. For example, in the visible spectrum their properties can range from

transparent6 to completely opaque to highly reflective.

7 Structurally, sapphire (Al2O3) is among

the hardest materials known to man. Additionally, oxides are typically brittle in nature and yet

some have the incredible ability to change shape through expansion or contractions when an

electric field is applied (piezoelectric effect8). In terms of electronic properties, oxide materials

seem to run the gamut of achievable states. Although commonly thought of as insulating

materials, as is the case with HfO2 and ZrO2,9 oxides can also be metallic

10,11 or

semiconducting.6,12

Going one step further, some of the best known high Tc superconductors are

oxide based materials.13,14

Even the insulating oxide materials can have “hidden” electronic

properties such as piezoelectricity,8,15-17

two-dimensional electron gases,18-20

and

multiferroism.21-23

Colossal magnetoresistance, has been demonstrated in the manganite

family.24

Oxides get even more interesting when nominally forbidden properties are merged

such as polar metals25,26

or transparent conducting materials.27,28

Many of the electronic properties in oxide based materials are coupled with their crystal

structure and elemental constituents. For instance, piezoelectricity is inherently tied to the

crystal lattice.8 For this reason, perovskite based oxide materials are some of the most

5

interesting materials to study. The perovskite unit cell is shown in Figure 1.1 with A cations on

the corners, the B cation at the body center position, and the oxygen at the face centers. Although

classically a perovskite is strictly a cubic material, which has the chemical formula ABO3, the

definition is commonly relaxed to include many other crystallographic symmetries as well as

slight variations of components such as A1-xÀxBO3-δ. Even the parent structure, CaTiO3, was

originally believed to be cubic and later determined to be orthorhombic as the experimental

methods improved.29

The basic flexibility of combining two cations with three oxygen atoms

makes this material system extremely interesting to study as a way to tune electronic parameters

by exchanging one element for another. Goldschmidt et al. created a general rule by which one

can determine which elements can go together in a perovskite unit cell based on their relative

atomic radius back in 1926.30

Typically, the A element in the perovskite structure originates

from the Alkaline Earth series or Lanthanide row, while the B element is often from the

Transition Metal group. This results in a massive number of possible combinations, which is

Figure 1.1. The classic perovskite unit cell of CaTiO3 with green Ca atoms on the

corners, a blue Ti atom at the body center, and red oxygen atoms at the face center

positions. The oxygen octahedra surrounding the Ti atom is shown.

6

why perovskites are so interesting: one can tailor the chemical formula to access a huge range of

electronic, magnetic, and optical properties.

A lot of the interesting properties that exist in oxides are observed in the bulk crystalline

form but things get even more interesting when considering the epitaxial thin film form. In thin

film form, the “bag of tricks” transforms from adjusting the chemical composition to adding

many new ways of molding these materials by changing their electronic state. With thin films

we can access interfacial states, form metastable states far from equilibrium, create complex

artificially layered heterostructures or superlattices, and tune the strain. This is not an exhaustive

list, but the bulk material properties can be formed, altered, or enhanced creating a seemingly

endless possibility of combinations to achieve nearly any desired electronic state.31

Additionally,

thin films are far more practical for industrial applications as bulk materials are far more difficult

to incorporate in device structures. Clearly, oxide thin films have a lot of potential in the area of

electronic materials and have garnered a lot of attention over the past few decades.

1.3. Deposition techniques

Electronic oxide thin films are clearly an interesting discipline of material science to

study and a lot of theory has been created predicting many interesting phenomenon in this area.

However, actually forming epitaxial thin films and creating complex heterostructures is a field of

study in it of itself. Creating phase pure high quality single crystalline thin films is important for

probing the true intrinsic properties of the material and requires the film to be as close to defect

free as possible. Additionally, the substrate material and lattice mismatch concerns are important

when depositing films since the substrate can impart its physical properties into the film and acts

7

as a template upon which the film can be formed. For strain engineering in particular, the lattice

mismatch can be extremely important. Figure 1.2a shows some of the frequently used oxide

materials and their relative lattice parameter.32

This diagram is an easy reference for assessing

some of the possible material combinations and for choosing the appropriate substrate material

for the targeted strain state. A schematic example of compressive, tensile, and zero strain is

shown in Figure 1.2b, c, and d where the film adopts the in-plane lattice parameter of the

Figure 1.2. Common lattice parameters and strain states. a Shows film and substrate lattice

parameters and is often referred to for relative strain state identification.(adapted from30

) In b-d,

compressive, unstrained, and tensile strain states are demonstrated respectively. The in-plane lattice

of the bottom substrate unit cell is them adapted by the film above.

8

substrate, known as coherent growth. By choosing a substrate and comparing its lattice

parameter to the bulk value of the film material, particular strain states can be engineered if the

film is able to grow coherently. However, if the strain state is too large the film will relax and

assume the bulk value or the substrate will not act as an adequate template and single crystal

growth may become unlikely.

Strain engineering and substrate selection is only one of the many free variables that need

to be considered when growing thin films. This is what makes epitaxial film deposition a bit of

an art form; with so many varying parameters, it is often impossible to scan the entirety of the

parameter space. As such a sense of intuition and a systematic approach may be required to form

the desired phase. What parameters that are controllable are dependent on the growth technique.

For the purpose of this thesis, we will only introduce MBE, PLD, and sputter deposition. Other

techniques exist but these are three of the more common growth techniques for complex oxide

deposition. MBE will be introduced here and referenced later in the thesis but all the films

presented here were primarily grown by sputtering and occasionally by PLD.

1.3.1. Molecular beam epitaxy (MBE)

Molecular beam epitaxy (MBE) is a thermalized process by which constituent materials

are evaporated in ultrahigh vacuum and directed onto a substrate. While the evaporation of

material as a thin film deposition technique has been around for over a hundred years,33

the use

of evaporated species to form single crystalline materials was really developed at Bell Labs in

the mid 1970’s.34

In this work, they developed the ability to control the fluxes to account for the

changes in sticking coefficients of different species in order to grow GaAs and related

9

superlattices. The vapor species impinge on the heated substrate with energies typically less than

1 eV.35,36

This low energy makes this growth technique very gentle in order to avoid a lot of

damage. Furthermore, the temperature of the heater can be tuned to adjust the sticking

coefficient but also provide further thermal energy to allow for surface diffusion. This process is

rather slow and requires an extremely clean system and ultrahigh vacuum in order to avoid

contaminants. MBE is capable of producing some of the highest quality thin films and offers

good stoichiometric control. A schematic of an oxide MBE system is shown in Figure 1.3 for

reference.37

Typically with MBE, to achieve such a high degree of stoichiometric control, some form

of in situ calibration is required. Often included in an MBE system is a quartz crystal monitor,

which uses the changing of frequency to determine the mass change. Thereby the rate of

Figure 1.3. A schematic of an oxide-MBE chamber is shown in a with common components including

a quartz crystal monitor and a RHEED system. In b a zoomed in view is provided demonstrating the

ability to shutter one source while depositing from another to achieve atomic layered deposition.

(adapted from35

)

10

deposition can be determined. In addition to the quartz crystal monitor which cannot be used

during deposition, RHEED is also commonly used. RHEED offers rea-time diffraction

information of the sample surface which is particularly useful for monitoring layer by layer

growth. A much more detailed introduction to RHEED will be given in Section 1.4.

Oxide MBE came about in the late 1980’s38

with the onset of cuprate based high Tc

superconductors. These types of superconductors require high partial pressures of oxygen,

which is counterintuitive to MBE growth since MBE requires ultrahigh vacuum. To avoid this

issue, groups developed oxygen plasma and ozone based oxygen sources which are leaked into

the chamber directly in front of the sample. Not only does this keep the total pressure in the

chamber very low but it also takes advantage of highly reactive oxygen species. Due to this low

total pressure, there was no issue with incorporating RHEED into these systems as there was not

enough oxygen present to damage the filament used in the electron sources. The tuning

parameters during oxide MBE growth are somewhat straightforward as one can tune the

substrate temperature, the ozone pressure, and the source temperature or evaporation rate. This

is not to say that MBE growth is trivial by any means but is closely governed by thermal

dynamics. Oxide MBE continues to be used in a similar fashion in order to produce some of the

high quality materials at a predominately research level.

1.3.2. Pulsed laser deposition (PLD)

While MBE and pulsed laser deposition (PLD) are both used to grow epitaxial thin fims,

materials they have some major differences. First, PLD was originally established as a growth

technique specifically for oxide high Tc superconductor deposition in the late 1980’s,39

unlike

11

MBE which began in the era of III-V semiconductors in the 1970’s.34

In fact, PLD is not

efficient in depositing from metal sources due to their high degree of reflectivity in the UV

wavelength and is truly optimized for ablating ceramic materials that have a high degree of

absorption of UV light. PLD uses a high powered excimer laser (248 nm) to ablate a target

material which results in ablated species energies on the order of several hundred eV.35,36

This is

vastly different from the less than 1 eV energies observed in MBE. These types of kinetic

energies associated with PLD can result in damage to the substrate but also gives the impinging

atoms a lot more energy to move around on the surface of the substrate to form epitaxial layers.

To a certain degree the impinging energy can be tuned through partial pressures in the chamber

and working distances. Higher pressures in the chamber with larger working distances results in

more collisions before reaching the substrate and in this way a lot of energy can be removed

from the atoms. While PLD can operate in nearly any pressure, the higher degree of vacuum that

is used, the more likely damage is to occur. The pressure in the chamber must also be tuned for

the desired oxidizing state. Typically, molecular oxygen is used to control the pressure and is

not particularly reactive but can be used to avoid a reducing environment which some

compounds are extremely sensitive to (e.g., SrTiO3 (STO)).

One of the main benefits of PLD is the high degree of stoichiometric transfer from the

target material to the substrate. In this way, a ceramic target can be formed by mixing the

desired chemical components in exactly the desired ratio which will subsequently be transferred

to the substrate. This process, however, is tied to the laser energy and therefor the stoichiometry

can be tuned to a certain degree through the laser energy.40,41

While ablating from a

stoichiometric target has its benefits, the same degree of control that has been shown in MBE in

order to form higher order phases such as Ruddlesden-Popper phases, has not been as

12

ubiquitously shown in PLD. Only more recently has the ability of PLD to have fine

stoichiometric control from two binary oxide targets compared to a single target been

demonstrated.4

Despite PLD having a level of stoichiometric transfer, forming high quality epitaxial

films is not necessarily straight forward in large part due to the high energy of the process.

There are many parameters to tune in order to obtain the optimal growth position. As previously

discussed, the pressure and working distance in the chamber can be important parameters for

tuning the energy but are not independent parameters and, as such, cannot simply be tuned

linearly in order to be optimized. In addition to the pressure and working distance, there is the

substrate temperature and all of the laser parameters. The laser parameters include laser energy,

magnification level, and mask size. These are important parameters and are not necessarily

straightforward in terms of how they impact the growth. The dependency of all the free

parameters on one another is a large part of what makes PLD a complicated growth technique.

None the less, PLD has emerged as one of the widely used research tools for depositing high

quality epitaxial oxide films.

Similar to MBE, PLD also takes advantage of in situ RHEED as a powerful real-time

analysis tool which can greatly assist in the optimization of the film growth and provide unit cell

control. In PLD however, further care needs to be taken due to the potential for high pressures of

oxygen in the system which can easily oxidize the filament in the electron gun and break it. To

avoid this process Rjinders et al. established the high pressure RHEED system which uses a

small aperture and a double differentially pumped extension tube to maintain 10-5

Torr or better

in the filament cavity.42

Further discussion and description of this will be provided later in the

thesis in Section 1.4. Since the development of the high pressure RHEED system, RHEED has

13

become a staple in PLD chambers and greatly assists in the deposition of complex

heterostructures.

1.3.3. Sputter deposition

Unlike the previous two deposition techniques, MBE and PLD, sputter deposition is

firmly rooted in semiconductor industry as a metallization technique.43

In addition, sputtering

has been around for the longest amount of time having been discovered in the 1850’s by W. R.

Grove.44

Incidentally, this was an accidental discovery while he was studying dc glow discharge

tubes which are essentially analogous to the sputtering process. Essentially, a low partial

pressure of a particular gas is isolated in a vacuum cell with an anode and cathode. By applying

a large enough bias between the two electrodes, the gas will ionize and create a plasma. This is

in effect a characteristic glow which can be used in spectroscopy experiments or simply as a light

Figure 1.4. Example of sputter processes. a and b compare a conventional sputter source with a

magnetron sputter gun. The magnetic field contains the plasma close to the source by trapping the

electrons and causing an increase in argon ions near to the target.(adapted from43

) An example from one

of the early high Tc superconductor sputter depositions in c, showing the 90° off-axis sputtering geometry

used to avoid substrate damage from oxygen ions.(adapted from44

)

14

source in the form of fluorescent lights used every day. What Grove discovered was that with

time, the cathode material disappears. In particular, when he varied the gas species from a

hydrogen environment to one rich in nitrogen the cathode surface went from a polished surface

with hydrogen gas to one with a hole forming in the nitrogen environment. Essentially, without

directly saying it, Grove found that heavier elements sputter more efficiently than lighter

elements. While he was not directly looking to identify a thin film deposition technique, his

observation spurred a new field of study that is still relevant nearly 200 years later.

What Grove essentially discovered is the sputtering process which is in effect the process

that occurs when ionized gas is accelerated toward a target material. This bombardment of

atoms on the target surface can cause the removal or “sputtering” of the target material. See

Figure 1.4 for a schematic of sputtering processes.45,46

The sputtered species will then condense

on any surface that the material can reach. In this way, thin films composed of the target

material can be formed. This process is fundamentally dependent on the gas species used, gas

pressure, and the accelerating potential. Sputtering is a momentum transfer process and, as such,

the higher the energy and higher mass of the impinging sputter species the higher the deposition

rate will be. This is what Grove observed as well, with hydrogen being an extremely inefficient

sputter species due to its low mass.44

Typically, argon is used as a sputter species due to its

modest mass and also the fact that it is a noble gas means that it will not react with the target

material in an undesirable way. The gas pressure is an important parameter as the sputter

process relies on ionized gas species to exist. The ionization process is a result of gas atoms

colliding with free electrons and one another in the vacuum in order to knock out an electron and

begin a cascade effect. If the pressure is too low, there will not be enough collisions to begin the

cascade process forming a plasma and if the pressure is too high, the mean free path will be

15

insufficient for the atoms to gain enough energy between collisions to knock out electrons.

Furthermore, the pressure can have an impact on the sputter rate and the deposition rate which

are both dependent on the mean free path of the ionized gas and sputtered material, respectively.

The discovery of magnetron sputtering in the late 1960’s to early 1970’s enabled a more

efficient sputter process and also expanded the pressure range which sputtering can occur at to

lower values.47

The magnetron sputter source is included in the schematic shown in Figure 1.4

including the magnetic field lines that it produces and the outer annular and inner cylindrical

permanent magnets. The field lines act to contain the free electrons close to the cathode and aide

in the ionization process, particularly at low pressures. Magnetron sputtering quickly became the

norm for sputter deposition and is by far the most common sputter technique seen today.

However, for the purpose of this thesis and for in situ analysis techniques using electrons, the

presence of these magnetics is severely detrimental. Chapter 5 is dedicated to solving this issue

for the use of RHEED with sputtering in real time.

Sputter grown epitaxial oxide thin film deposition was demonstrated around the same

time as for PLD and MBE in the late 1980’s with the discovery of high Tc cuprate

superconductors.46,48,49

It was demonstrated at this time that 90° off-axis sputtering is an

important growth geometry for oxide deposition as high energy oxygen ions formed in the

plasma can be accelerated opposite the cathode. If the substrate is placed on an anode directly

opposite the cathode then the oxygen ions can actually act to sputter the substrate material and

cause significant damage. To avoid the negative oxygen ion bombardment, the substrate is

placed 90° with respect to the sputter gun surface as shown in Figure 1.4.46

Additionally, a high

pressure of gas is often used to increase the number of collisions before reaching the substrate so

as to create a more “gentile” growth environment. In general, sputtered species can have a very

16

large range of energies from 1 eV to hundreds of eV depending on these growth conditions,50,51

which makes it an adaptable and interesting growth technique compared to MBE or PLD.

What makes sputtering a particularly interesting growth process in comparison to PLD

and MBE is the active material species it can create. For instance, in oxide MBE, a highly

reactive ozone source is often employed to provide oxygen to the substrate and ensure a fully

oxidized state, or alternately molecular oxygen in PLD (far less active than ozone or atomic

oxygen). Whereas in sputtering, active atomic oxygen can easily be formed from the oxide

target or by introducing molecular oxygen and for applications such as high Tc superconductors,

this can be extremely advantageous as higher oxidized states are often required. Similarly, in

PLD or MBE growth of nitrides, a nitrogen plasma is often incorporated into the system to crack

the strong triple bond that exists in N2 and create a stoichiometric film. In sputtering, the plasma

inherently exists and, as such, no additional process is needed other than flowing nitrogen gas.

This is true for many material systems such as nitrides, oxides, carbides, etc. where a gas species

is incorporated into a film. Additionally, sputtering is exceptional at depositing metals and in

particular refractory metals (and alloys), which can be very challenging or impossible to grow

using PLD or MBE. This is due to the fact that sputtering is a purely kinematic process based on

momentum transfer. Because of this breadth of possible material systems deposited by

sputtering, it is a highly versatile growth process that is likely held back by its lack of in situ

analysis options.

1.4. Reflection high energy electron diffraction (RHEED)

17

Nishikawa and Kikuchi conducted the first reflection high energy electron (RHEED)

experiment in 1928 as an alternative to transmission electron diffraction.52

Since then, RHEED

has been developed into one of the principal surface structural analysis techniques particularly as

a monitoring technique coupled with epitaxial thin film growth. Many works have been

published regarding this topic but we have found Ichimiya and Cohen,53

and also Braun’s54

books regarding RHEED extremely useful and this section draws heavily from their work.

While RHEED can be used independently of any deposition process, it is most powerful

when used in combination with an epitaxial growth process. This is due to the small penetration

depth and low incidence angle leading to strong surface sensitivity that is hard to achieve with

other in situ techniques. Additionally, since RHEED is a grazing incidence technique, it is easily

coupled with a growth process or other analysis technique since the bulk of the physical space in

front of the sample is completely unoccupied. While RHEED can provide information regarding

nearly any surface state from amorphous and rough to polycrystalline and faceted, it is most

useful in conjunction with epitaxial growth and atomically smooth surfaces since the most

detailed surface diffraction information can be extracted from patterns stemming from these

surfaces.

RHEED is also surprisingly easy to employ at the most basic level consisting only of an

electron gun and phosphor screen. Naturally, as the desired functionality and application get

more complicated, so does the RHEED apparatus. Even in some of the more complicated forms,

the basic principle and setups are relatively simple. Similarly, the RHEED analysis can be very

straightforward and simple to understand via eye inspection at the fundamental level with

increasing complexity as the desired knowledge gets more complex at times requiring detailed

18

video capturing systems and analysis software. This thesis will not dive into the deep underlying

principles associated with RHEED and will focus on more basic concepts.

As the name describes, RHEED is a form of electron diffraction that occurs resulting

from a reflection off of a solid surface at energies in the 5 – 100 keV range but more commonly

20 – 35 keV. The higher energy electrons are required to expand the Ewald’s sphere to the point

that multiple reflections are intersected simultaneously and, as such, more than one diffracted

spots are observable on the phosphor screen. For amorphous solids, there are naturally no

diffraction spots and only a diffuse background intensity is observed. For polycrystalline

materials, the reciprocal lattice is transformed into spheres and its cross section with the Ewald’s

sphere is a circle. As a result of the shadowing process caused by the sample at low incident

angles, a set of incomplete circles is viewed on the phosphor screen stemming from the various

scattering planes. When the surface becomes very rough and 3-dimensional, yet crystalline in

nature, a set of repeating spots will appear on the screen similar to a grid. This results from the

electron beam penetration depth being larger than the 3D island sizes and, as such, multiple

scattering events occur with different scattering conditions. This diffraction pattern is then

unaffected by the incident and azimuthal angles.

For high quality single crystalline materials with smooth 2-dimensional surfaces, the

diffraction pattern becomes far more interesting and a lot more detailed information can be

obtained. For these materials, the diffraction pattern is extremely sensitive to both the incident

angle of the electrons on the sample surface and the electron beams relative alignment to the in-

plane crystallographic orientation. As such, control over these two parameters known as the

“tilt” angle and the “azimuthal” angle is extremely important. For single crystalline diffraction,

the finite reciprocal lattice spots intersect with the relatively large Ewald’s sphere to form

19

distinct diffraction spots on the phosphor screen. The spots fall onto a ring around the

transmitted spot known as a Laue zone. Multiple Laue zones can be viewed and are a function

of the wavelength of the electrons, with each zone occurring at integer multiples of the

wavelength. As the wavelength of energy is changed, the number of diffracted spots visible on

the screen for a given incident angle will change. As the energy increases, more spots will be

observed as the pattern contracts. These spots can be indexed based on the azimuthal angle but

the central spot perpendicular to the sample surface is always the (00) specular spot for a 001

oriented crystal. In the [100] azimuthal direction, the next set of symmetric spots around the

specular spot would be referred to as the (10) and (1̅0).

The spacing between the diffracted spots is indeed a function of the plane spacing in the

crystal. Since RHEED is so surface sensitive, this spacing can be monitored to follow the top

most layer’s lattice parameter as a function of time or thickness. This may be particularly

interesting for detecting changes, such as lattice relaxation during growth. One can immediately

determine at what thickness the film relaxes back to its bulk state and know the limit for coherent

growth.

By observing the RHEED pattern, if the crystal surface is good enough, it can be

understood very quickly if there is a change in the crystal symmetry through the appearance of

half order peaks, or something similar, based on the crystal symmetry and orientation. Similarly,

and commonly seen with Si and other classic semiconductors, one can observe if there is any

surface reconstruction taking place. This would typically take the form of the series of evenly

spaced spots between the main diffraction spots that correspond to the reconstruction order. For

instance, in a Si (111) oriented crystal, a 7x7 surface reconstruction is known to exist and this

can easily be seen with RHEED as 6 distinct spots will appear between the main diffraction

20

spots, with one diffraction spot counting towards the reconstruction order. This demonstrates

how surface sensitive RHEED can be and how useful it can be to know this information in real

time before or during a deposition.

Arguably the most commonly observed and desired phenomenon when discussing in

situ RHEED is the intensity oscillations attributed to a mono layer or single unit cell of

deposition. This gives an unprecedented real time growth rate calibration as well as precise

atomic level control of the growth. Intensity oscillations were first discovered during MBE

growth of GaAs in 1981 by Harris et al.55

and have since become a key feature to most epitaxial

growth articles using PLD or MBE. Figure 1.5, adapted from Ohring’s56

book, provides a nice

depiction of the origin of RHEED intensity oscillations although in reality it is a more

complicated process than described here. A good example of this is that the oscillation phase

Figure 1.5. Schematic representation of RHEED intensity oscillations. The

RHEED intensity is shown to vary as a function of layer coverage. (adapted from54

)

21

can shift as a function of incidence angle of the electrons which suggests, that the peak in

intensity does not always indicate a completed surface.57

This can depend on your diffraction

conditions, but the period of oscillations is directly related to the time needed to deposit one unit

cell.

The ability to observe RHEED intensity oscillations during sputter deposition is the

ultimate goal of this thesis as it had not previously been observed. Short of using synchrotron x-

ray source, very little is known about the growth mode during sputter deposition. As previously

discussed, sputtering is a physical vapor deposition like PLD but its energetics and active species

can be completely different from PLD. Additionally, sputtering is not a thermal process like

MBE and it falls into an unknown parameter space in terms of growth dynamics. This makes the

combination of RHEED with sputtering an extremely interesting and exciting area to pursue.

1.5. Outline of thesis

Chapter 1 is designed to provide motivation for why pushing the bounds of in situ analysis

during oxide thin film deposition is an important area to study. In particular for epitaxial thin

film deposition by sputtering where relatively few real time analysis options are available. The

importance of oxide thin films and breadth of properties is described. A brief comparative

summary of molecular beam epitaxy, pulsed laser deposition, and sputter deposition is given

with an emphasis on oxide growth, the strengths and weaknesses of each, and the availability of

in situ RHEED. A history of RHEED is given as well as a description of how RHEED works.

Several of the important characterization methods by RHEED are outlined and explained.

22

Chapter 2 gives a detailed description of a novel “hybrid” deposition chamber that we designed

and constructed. This chamber combines sputter deposition, PLD, and RHEED into one simple

system for the use of depositing dissimilar materials such as metal-oxide heterostructures. The

chamber and control software written in LabVIEW are described in detail. As a way to

showcase the power of this system, a capacitive heterostructure for use in a superconducting

quantum bit was deposited and the microwave loss properties are shown.

Chapter 3 is intended to demonstrate the growth of a scientifically relevant material system,

LaAlO3/SrTiO3, by sputter deposition. This acts in two fold; first to demonstrate the interesting

properties that exist via a scalable growth process, and second to show the strong thickness

dependence on the electronic properties as an ideal example of a system where RHEED would

be useful. The LAO/STO thin films were grown in a pure argon environment as a way to control

the limit the partial pressure of oxygen which is critical for LAO growth. Additionally similar

electronic properties were seen compared with PLD growth LAO films and in particular the 4

unit cell critical thickness. And finally, conductive nano-wires in the films were shown to be

writeable and erasable via conductive AFM demonstrating their device applicability.

Chapter 4 demonstrates for the first time RHEED oscillations during sputter deposition starting

with the LAO/STO heterostructure. These films are grown with a single gun in the 90° off-axis

geometry and although the RHEED pattern is tilted and distorted the specular spot intensity

could still be monitored. The challenges of combining RHEED with sputtering are clearly laid

out in this chapter. Additionally, oscillations were seen in the LSMO/STO and SRO/STO

23

heterostructures demonstrating the robustness of this technique. Both the LSMO and SRO films

exhibit significantly stronger RHEED oscillations than LAO and interestingly the SRO shows

oscillations out to greater than 50 periods which has never been reported before for PLD grown

films.

Chapter 5 is dedicated to better understanding the difficulties of incorporating RHEED with

magnetic fields produced by magnetron sputter sources. Finite element models were constructed

for various growth geometries and in particular single gun and dual gun 90° off-axis

arrangements. Using these models we were able to not only predict the degree of deflection of

the electron beam but were also able to find a way to mitigate the deflection to the point that the

beam travels nearly straight. In this way we are able to experimentally show a significantly

sharper RHEED pattern in the presence of two “anti-symmetric” sputter sources with minimal

deflection. We again demonstrate that extended RHEED oscillations are observable in this setup

during SRO growth on STO. This arrangement is extended further to be used as a general rule of

thumb for many growth geometries from 90° off-axis to purely on-axis.

Chapter 6 is intended to summarize this work and draw all the chapters together. Additionally, it

is aimed at giving insight into the further progression of RHEED with sputtering including

thoughts on interesting areas to pursue and in particular material spaces that are either impossible

or very difficult to axis via other growth techniques.

24

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31

2. PLD-Sputtering Hybrid with RHEED

2.1. Introduction

As previously discussed, high crystalline quality epitaxial structures have an incredibly

broad range of applications and very interesting properties can be engineered and studied using

the thin film approach. The epitaxial deposition of dissimilar materials (e.g., metals on oxides) is

promising for many device architectures. Theoretical predications indicate electric field control

of magnetism is possible at interfaces between magnetic metal and dielectric thin films.1

Furthermore, the deterministic switching of a ferromagnet at room temperature by an applied

electric field to the underlying multiferroic has been recently demonstrated.2 This chapter will

focus on an epitaxial shunt capacitor component for solid state quantum computing devices

(SSQCD) consisting of superconducting metal electrodes and dielectric interlayer3,4

as a way to

showcase the powerful custom built “hybrid” deposition system specifically designed to tackle

these challenging structures.

A limitation in SSQCD devices is the loss of information that originates from dielectric

loss (tanδ).5,6

It has been shown that the energy relaxation rate is dominated by spurious coupling

of the qubit to low-energy two-level state (TLS) defects in the amorphous dielectrics of the qubit

circuit.5 These defects are believed to arise from atomic scale structural imperfections

7 and are

known to exist in the surface oxides of the superconductors, at the superconductor-insulator

interface, and in the bulk of the amorphous dielectrics of the circuit.8,9

Typically, SSQCD

32

structures employ amorphous or polycrystalline dielectric thin film layers, where grain

boundaries, poor interface quality, and other microstructural defects can be a significant loss

mechanism.10

Single-crystals are difficult to prepare for electrical measurements and are not

applicable to high-density device architectures required for SSQCDs. Therefore, the use of

epitaxy provides an opportunity to improve loss characteristics and the ultimate performance of

SSQCDs.

Combining metal layers with oxide layers during heteroepitaxial thin film growth creates

several critical issues that need to be addressed. First is the concept of choosing the appropriate

growth technique. As was addressed in Chapter 1, different techniques whether it be sputtering,

PLD, or MBE, perform well for different material systems. For the proposed material system in

this section of a heterostructure containing both oxides and metals, care needs to be given to

which deposition technique is employed. PLD is a proven technique for depositing high quality

epitaxial oxide dielectrics with superb stoichiometric transfer and flexibility (e.g., multiple

targets).11,12

However, PLD is inefficient for depositing materials with high thermal

conductivities and high optical reflectance in UV range, such as metals.13

As such, incorporating

metals into this stack is problematic when using only PLD and therefore another grow method

should be used. For metal growth and in particular refractory metals, magnetron sputtering is

particularly well suited. For this reason there is a need to implement both sputtering and PLD for

the deposition of this type of film stack.

As previously mentioned, the material stack for the SSQCDs is very sensitive to the

interface quality. With this in mind, it is critical to deposit the entire stack in one vacuum system

to avoid contaminating the interface by exposing it to air when changing from a sputter chamber

to a PLD chamber. Exposing the sample to air can cause carbon and hydroxyl impurities among

33

other things to form on the surface. While many oxide materials are extremely stable in a wide

range of environments, a perfect example of these impurities is demonstrated in Figure 2.1 where

an AFM image of the well-known dielectric MgO is shown immediately after being exposed to

air and then again several hours later. MgO is a hygroscopic material,14

meaning that it quickly

absorbs water and forms a hydroxyl which is the effect that is being observed on the surface in

Figure 2.1b. To avoid this type of contamination it is essential to keep the sample in good

vacuum until a stable capping layer is deposited.

One also needs to consider the metal layer when designing the proposed capacitor

structure. Some metals are incredibly inert, such as Pt and Au, and do not readily form impurity

phases whereas metals like Ni or Al very quickly oxidize at room temperature in air. Many

metals fall somewhere in between and form oxides at higher temperatures. This is another

reason for depositing the entire trilayer in situ without breaking vacuum so as to avoid any

oxidation of the metal layer.

Figure 2.1. MgO surface after exposure to air. AFM images from an MgO thin film

taken immediately after removing from the chamber in a and several hours later in b.

The surface quickly degrades when exposed to air.

34

This requires one to also consider that the proposed capacitor stack involves a metal layer

sandwiched between two oxide layers. If considering a polycrystalline stack with metal and

dielectric layers, they could easily be formed at room temperature with a greatly reduced risk of

oxidation of the metal layer. However, in this work the goal to make the dielectric layer single

crystalline which requires the metal layer below it to also be single crystalline since it acts as the

template layer. Not only does this limit the material choices to ones with good lattice match, but

it also limits the metal layer to one that does not oxidize easily since high temperatures are

typically needed to form single crystal phases. Any parasitic oxide that forms at either of the

metal-oxide interfaces would greatly increase the energy loss in the system. This also

necessitates that the growth chamber be able to change environments efficiently from a pure Ar

environment for the sputtering of metals to one which contains oxygen gas for the oxide growth

and back to the metal environment without any oxygen contaminates left behind.

During a complex multistep growth process there are a lot of variables which can

significantly influence the resulting properties of the sample. To successfully fabricate this

challenging SSQCD epitaxial stack, the use of an in situ analysis method is vital. RHEED is an

excellent candidate as an in situ monitor due to the vast breadth of information it can provide.

Of particular importance is observing and understanding the interface state during and between

layers. RHEED is a diffraction technique that is very sensitive to the top few layers . As such,

RHEED patterns give information to the surface smoothness (2D or 3D surface) and also to the

crystalline quality (polycrystalline, amorphous, or single crystal). All this information can be

determined quickly via visual inspection of the pattern. Information such as relative lattice

parameters, film relaxation, and thickness oscillations can also be determined but requires further

analysis.

35

The remainder of this chapter will be dedicated to describing in further detail the custom

hybrid PLD-sputtering chamber that was built in order to establish the ability to grow dissimilar

materials in one chamber to preserve the interface quality and efficiently grow multiple layers.

This chamber was equipped with a RHEED system, which can be used during either growth

technique via a simple 180° rotation of the heater. Additionally, the custom LabVIEW program

used to control the hybrid system will be outlined and described. The latter part of this section

will be devoted to the growth and characterization of several epitaxial shunt capacitor structures

that were made using this system.

2.2. System design

2.2.1. Deposition chamber

A combination of PLD and sputtering is a desirable solution for the fabrication of

multilayers of dielectric and metal epitaxial thin films, exploiting the strengths of both growth

techniques. A critical aspect of multi-source deposition systems is to avoid exposing the sample

to air, which causes interface contamination including carbon and hydroxyl impurities.

Therefore, multiple sources are either integrated into a single chamber or a cluster style system.

Cluster style systems typically require multiple chambers, many sets of pumps, duplicates of

gauges, manipulators to move the sample around, etc. This lends to cluster styles being both

expensive and complex with many moving parts resulting in long times scales to switch between

sputtering and PLD.

36

On the other hand, a typical single vacuum chamber containing multiple growth sources

lacks a certain degree of flexibility. All growth equipment must be placed in a hemisphere

around the fixed heater, limiting the space and geometries available. The PLD and sputter

targets cannot both be placed in an on-axis configuration because they would both occupy the

Figure 2.2. Schematic layout of the "hybrid" deposition chamber showing the main

components. Specifically the opposite facing sputter sources and PLD targets. The heater can

rotate between the two growth geometries and obtain RHEED patterns during both depositions.

37

same physical space. The RHEED geometry also needs consideration since it requires precise

alignment with the sample at a grazing incidence angle; it also limits the space which the growth

sources can occupy.

To overcome these limitations while also maintaining a simple, robust, and relatively

inexpensive single chamber setup, a “mirrored” geometry was developed by which the PLD

growth occupies one hemisphere of the chamber and a simple 180° rotation of the heater

accesses the other hemisphere containing the sputter geometry. Figure 2.2 shows the chamber

design and demonstrates this mirrored design.

Figure 2.3. The mirrored growth geometries are shown more clearly in a and c.

The RHEED patterns from an Al2O3 substrate are shown in b and d before and

after the 180° rotation between the sources.

38

The RHEED setup also benefits from the mirrored geometry. Since the heater rotates +/-

180° around the central axis, which is also approximately in the same plane as the sample

surface, only a small adjustment to the incident angle of the electron beam is required to

reestablish a diffraction pattern after rotating between deposition sources. The same diffraction

plane exists after the rotation and only the plane normal of the substrate has changed. This is

demonstrated in Figure 2.3 where a mirrored c-plane Al2O3 substrate diffraction pattern is shown

in both the PLD and sputtering heater orientations. This enables the easy integration of in situ

RHEED monitoring during both PLD and sputtering growth with only a several second lapse

while changing between alignments. To our knowledge, no other single chamber multi-source

design has incorporated this unique and practical approach.

While a certain level of sensitivity in sample alignment exists for PLD and sputter

deposition, the degree to which they are sensitive is far less significant when compared to the

RHEED alignment. RHEED requires very precise sample positioning in order to achieve desired

diffraction conditions that are reproducible. In the hybrid system there are five degrees of

freedom available for the sample positioning. The heater stage is mounted on an XYZ

manipulation stage with ±1-inch translation in X and Y and ±2 inches in Z.

There are also two angular rotations, tilt (Φ, ±180°) and azimuth (χ±90°). The azimuthal

rotation is used to adjust the in-plane angle of the sample with respect to RHEED electron beam

and has no function with regards to growth alignment. The tilt rotation serves two purposes;

moving between the PLD and sputtering growth hemispheres via the 180° rotation, and adjusting

the incidence angle (α) of the electron beam on the sample. This latter adjustment is the motion

that requires the most precision and to achieve optimal RHEED alignment a 1:70 gear reducer

must be combined with the stepper motor to improve the resolution to the hundredths of a degree

39

range. The angular motions and their importance for RHEED are well known and a thorough

description of them and their significance has previously been reported.15

All five of the degrees

of freedom are controllable through a custom built LabVIEW program. A more detailed

description of the LabVIEW program is given later in this chapter.

The heater itself is a resistive style heater with a current controlled dc power supply that

is rated up to 950°C in an oxidizing environment. The power supply is regulated through a

feedback loop with the temperature controller via a variable dc-voltage output. This type of

control is preferred over a relay because it prevents step functions in the output power which can

affect the trajectory of the electron beam adding undesired noise to the RHEED signal.

As previously touched on, the flexibility in this system is one of its major highlights. For

certain applications two on-axis sputter guns are desirable and as such this chamber has the

ability to incorporate both PLD as well as two sputter sources. In particular, for the SSQCD

growth a 1.33-inch magnetron sputter source for the rhenium (Re) target and a 2-inch magnetron

sputter source for the aluminum (Al) target were mounted opposite the PLD target stage, as seen

in Figure 2.2, in order to form the two superconducting metal layers in situ. These sputter guns

were powered by a dc power supply. These guns were mounted on the conflat flange using

quick coupling flanges such that their working distances could be adjusted. Additionally, the

guns interfering with one another and collisions with the heater stage when it is rotating between

sputtering and PLD could also be avoided.

Opposite the sputter sources, the PLD target stage was mounted, which has the ability to

change between four materials via a four target carousel, again reaffirming the chambers

flexibility. The target stage is mounted on a linear feedthrough to maintain an optimal working

40

distance (~5 cm) but also avoid colliding with the heater stage while rotating. When changing to

the sputter geometry, the target carousel was retracted 3-5 cm. In order to maintain the desired

laser spot focus on the target, fine working distance adjustments were made by adjusting the

XYZ manipulator of the heater stage rather than adjusting the target stage. The laser spot has a

minimum size of ~0.03 times the original mask size with the current optics setup.

A base pressure of 2x10-8

Torr is achieved in the chamber without a bake out and

numerous Viton o-rings used. This pressure is achieved by the combination of a 520 l/s

turbomolecular pump and a 1,500 l/s cryo-pump. The cryo-pump is instrumental in not only

quickly achieving a good base pressure but also for rapidly changing growth environments (e.g.,

pumping out oxygen used for the dielectric layer and ensuring a nearly oxygen free condition for

the metal layer). The working pressure in the range of 1 mTorr – 1 Torr is set and adjusted in the

LabVIEW program and controlled with a throttle valve between the chamber and the

turbomolecular pump. The inlet of gas is controlled by a set of mass flow controllers for each

gas type. Once the flow rate is set there is a feedback loop between the motorized throttle valve

and the capacitance manometer gauge to maintain a very precise pressure in the chamber

regardless of the flow rate.

A load lock was later added to the hybrid chamber intended to limit the breaking of

vacuum and preserve air sensitive target materials such as the Alkaline earth oxide PLD targets

which are hydroscopic and decompose in air rapidly. The load lock is capable of the exchanging

heater blocks while maintaining high vacuum, but is incapable of transferring targets in and out.

This process requires the breaking of vacuum and removing the target carousel through the main

viewport door. The load lock is equipped with its own turbo pump, which spins up and down in

41

approximately 10 minutes making the transfer process fairly rapid. The load lock vacuum

components are also controlled through LabVIEW.

The high pressures (mTorr range) typically used during PLD and sputter deposition

necessitates the use of a double differentially pumped high pressure RHEED system16

to avoid

the oxidation of the filament in the electron source. The Mu-metal extension tube with a small

(~200 µm) aperture extends into the chamber to decrease the electron scattering due to the high

pressures. The aperture acts similar to a virtual leak with the main chamber at high pressure

relative to inside the tube, which is differentially pumped with a second turbo pump (<10-5

Torr

at the filament). The aperture has a very small conduction through it which is why a large

pressure differential can be achieved. The phosphor screen is symmetrically placed opposite the

end of the aperture with a total working distance of 20 cm. The entire electron gun arm is

attached to an XY manipulation stage to adjust for variable growth setups and heater positions

and also any magnetic deflection in the electron beam which will be addressed further later in

this chapter. The electron source and power supply are capable of emitting electrons with

energies ranging from 5 – 35 keV. Higher energies reduce the scattering cross-section and, as

such, in this work energies in the range of 30 – 35 keV were used.

Care and consideration needs to be paid to the combination of RHEED with magnetron

sputtering due to the relatively strong permanent magnets in the sputter sources. While this is

not the first demonstration of combining RHEED with magnetron sputtering,17

the reports of

doing so are rare, likely due to the significant bending of the electron beam as a result of the

permanent magnets located in close proximity. The rare earth NdFeB permanent magnets have a

surface magnetic fiel on the order of B ≈ 3,000 G and consist of an annular magnet around a

cylindrical magnet leading to the magnetic field not being straightforward to predict in three

42

dimensions. Modeling this magnetic field and its effect on the electron beam will be the subject

of an entire chapter (chapter 5) later in the thesis and is a major portion of this work. At this

stage it is sufficient to know the field exists and with the on-axis working distances for this work

the deflection is minor enough that the beam still strikes the 55 mm phosphor screen. A Mu-

metal enclosure was built for a single 1.33-inch magnetron sputter source, which reduced the

deflection further but was not compatible with multiple sputter guns. The dc-bias from the

sputter power supply is another possible source of electron beam deflection yet it was not found

to have a significant impact.

2.2.2. LabVIEW Control Program

A custom LabVIEW program, “Hybrid.vi”, was built to control the hybrid chamber. The

program has evolved through several iterations as needed, keeping up with the theme of

flexibility and robustness. As specific needs change, the program can easily be updated to

address these needs, which is not possible in most commercially designed software packages.

Initially the program had a built in function to capture and analyze the video taken of the

RHEED phosphor screen. However, the complexity of the required RHEED analysis outpaced

the ability to design the LabVIEW program and a commercial software, Kspace 400, was

purchased to handle the RHEED monitoring and analysis. All other chamber functions are

controlled through Hybrid.vi.

For gas handling and pressure control, Hybrid.vi interfaces with the mass flow

controllers, the throttle valve, the pneumatic gas valves, and pneumatic vacuum valves. The

pneumatic roughing valve and motorized gate valve for the turbo pump are controlled through

43

push buttons and light indicators to show their current states. Two gas types can be chosen at

one time with their respective scaling factor, automatically updating in the program. The

percentage of each gas type, the flow rate, and the total pressure can all be controlled in the

program through a set of drop down menus and numeric controls. A switch starts and stops the

gas flow and control process by fixing the set point pressure for the throttle valve, opening the

gas valves, and setting the mass flow controller flow rates. Once set, the pressure stabilizes

within seconds. The capacitance manometer and full range gauges are read into the program and

numerically displayed and the full range gauge is continuously updated in a plot of pressure vs.

time.

The five degrees of freedom of the heater stage are all controlled through Hybrid.vi by

individually moving each motor. Additionally, the heater can be adjusted by moving to one of

the 3 saved set points: Sputter, PLD, and Load/Unload. The drives automatically move to

specific positions where those events typically take place and then can be fine adjusted from

there. Another stepper motor is mounted to the target carousel and is used to select between one

of the 4 possible target positions and is controlled in Hybrid.vi. The heater temperature is also

set in the LabVIEW program and the actual temperature is plotted versus time.

A later addition to Hybrid.vi was the superlattice portion. In this section of the program

the excimer laser used for PLD can be controlled remotely. The repetition rate, voltage level,

and time or number of pulses can all be set and used to run the laser away from the main laser

control computer for individual layer growth. For growth of superlattices by PLD the laser

parameters, gas values, target position, and temperature can all be set independently for each

layer with adjustable wait times between changing parameters to allow for stabilization. The

desired number of laser pulses for each layer and the total number of layers can be set and the

44

program will run and reproducibly control each layer, removing human error and time. This is

especially useful when the repeat unit is reiterated 10s of times whether it is for superlattice

growth, sequential binary oxide growth,18

or any multi-target growth.

2.3. Metal – Oxide Heterostructures

In this section, two types of capacitive structures will be discussed: one with MgO as the

dielectric layer and the other with Al2O3 as the dielectric layer. We have grown many other

structures using various superconducting metals, including some nitride materials such as TiN,

and have also grown a broad range of dielectric materials (i.e. BeO, LaAlO3, LiF). The

challenge for growing any multilayer epitaxial heterostructure is finding materials with similar

lattice matches to limit the number of defects and to encourage well oriented growth. This

concept was introduced in chapter 1 along with a diagram demonstrating lattice mismatch.

The superconducting metal layer does not show a decrease in performance with single

crystal versus polycrystalline and as such the top metal layer does not need to be epitaxial. For

this work, polycrystalline Al was the material of choice for the top layer. For the bottom

electrode there are several critical considerations. The material should have good lattice match

to the substrate or in this case to the oxygen sublattice in the sapphire in order for it to grow

epitaxially. Since this layer will be sandwiched between two oxide layers it needs to be resistant

to oxidizing. Additionally, it needs to have a superconducting transition at an achievable

temperature and have low loss characteristics. For this layer, it was found that Re performs well

with respect to all of these criteria19-21

and will be the focus of the work presented here.

45

The Al/MgO/Re capacitor stack was grown on c-sapphire substrates which were annealed

in a flowing O2 gas environment at 1100 °C for 4 hours. This was done to improve the surface

quality of the as-received substrates prior to growth, which is critical for epitaxial growth. Once

the substrate is mounted on the heater block and inserted into the hybrid chamber, the multilayer

was deposited following the sequence shown in Figure 2.4 which depicts the temperature,

pressure, and growth technique as a function of time. The first layer of Re was deposited at 800

°C by sputtering in 20 mTorr of Ar flowing at a rate of 20 sccm. A dc power of 30 W was used

and the total growth time of 12 minutes and 30 seconds led to a Re thickness of 100 nm. A

working distance of 9 cm was used to keep the sputter rate high yet avoid back-sputtering of the

sample. Following the growth, a 2-hour anneal in vacuum was performed to help crystallize the

film further and improve surface roughness.

To better understand the growth dynamics, in situ RHEED was used to monitor the

growth of this heterostructure. Static RHEED images between layers and during the multilayer

growth can be seen in Figure 2.4 starting with the bare c-sapphire substrate. It is important to

note that Re grows with a 30° rotation with respect to the sapphire substrate orientation. The

relative layer growth orientations are shown in Figure 2.5 and are non-trivial. Once the Re

deposition starts, the RHEED pattern quickly changes, and by capturing a movie one can go back

and inspect the changes frame by frame. This demonstrates the true power of real time RHEED

as compared with a RHEED system that is separate from the growth system. It can be seen that

the original strong diffraction spots quickly weaken within 1 nm of growth, some indications of a

ring pattern appear and a diffuse background becomes noticeable. This indicates that during the

first few seconds of the Re deposition the surface of the sample goes from a sharp 2D surface

46

with high crystalline quality to showing signs of both polycrystalline nature as indicated by the

rings and some signal of an amorphous contribution which appears as a diffuse intensity.

Figure 2.4. Growth sequence static RHEED images. The general sequence of film growth during

the trilayer 100nm Al/50nm MgO/100nm Re on (0001) Al2O3 is shown in a with the partial pressures

and temperatures for each layer as a function of time mapped out. b-g show RHEED images at

different points throughout the growth starting with the bare substrate and ending with the MgO layer.

It can be seen that early in the Re growth the RHEED pattern is polycrystalline but soon recovers.

The MgO film starts with sharp spots but by the end has a more 3-dimensional surface.

47

The fact that there are some additional spots than the other features indicates that the film

is likely still very well oriented. It may also be a trace contribution from the substrate however,

if an electron mean free path through Re is assumed to be λe = 25nm,22

then the penetration depth

(D) is defined as D = λesin(α) where α is the incidence angle. The incidence angle is

approximately 1° which results in a penetration depth D ≈ 4.4 Å which is less than the total

thickness of the film taken at this point. This means that the spots are likely still coming from

the Re film and this becomes even more evident as the film gets thicker. After about 5 nm of

deposition, the rings and diffuse intensity significantly reduce and the spots become stronger and

more streak like. This is a well-known indication of smooth and two dimensional surfaces, more

similar to the starting substrate surface. By the end of the 100 nm deposition of Re, the spots

become even more well defined and all indication of a polycrystalline or amorphous film has

completely disappeared. This suggests that the growth may have a small reaction layer, but

rather quickly converts to a high quality epitaxial Re layer. This clearly demonstrates the power

of in situ RHEED as a way to understand the sometimes complex growth dynamics and

particularly during sputtering, which have rarely been demonstrated together.17

Following the Re growth the gas flow was turned off and the cryopump gate valve

opened in order to quickly evacuate the chamber of the Ar gas used for the Re growth. The

temperature was also adjusted to 450°C, which is the growth temperature used for the MgO

growth. Once the temperature was stabilized, the gas atmosphere was changed to 7.5 mTorr

oxygen with a flow rate of 20 sccm. The heater was also rotated 180° to the PLD growth

geometry and the RHEED pattern was realigned for the new incidence angle. The target

thickness for this layer was 50 nm with an approximate rate of 300 pulses/nm, so 15,000 pulses

were deposited from a ceramic MgO target.

48

At this temperature and pressure, the Re is stable and does not readily oxidize. This can

easily be confirmed by monitoring the RHEED pattern as a noticeable change would occur if the

film began to oxidize, which it did not. After 1 nm of deposition, the MgO RHEED pattern was

already noticeably changed although not to the same degree as seen during the Re deposition.

The spots and streaks remained throughout the MgO deposition but what is of note here is that

rather than the spots lying on a well-defined circle, the pattern becomes more rectangular and the

spots tend to lie on a straight line. This is typically an indication of 3-dimensional growth or a 3-

dimensional surface even though there are still well-defined spots. At this point the pattern tends

to look more like a transmission pattern with a regular array.23

However, MgO is cubic and is

being grown in the <111> direction on the hexagonal sapphire substrate and Re film. This can

simply be understood as a triangle matching on top of a hexagon, as seen in Figure 2.5, where

there are two possible orientations of the triangles, one rotated by 60° with respect to the other.

These are energetically equivalent and results in a twinning structure in the MgO film. This

twinning structure may contribute to the MgO RHEED pattern that was observed as there is a

doubling of the pattern. This doubling becomes more pronounced as the growth proceeds.

Immediately following the MgO growth, the heater is turned off and the chamber is

returned to high vacuum using both the cryopump and the turbo pump. The top electrode of the

multilayer stack is the Al layer which oxidizes very easily. To prevent this oxidation, it is

important to allow the heater to cool to room temperature (~20-50°C) and also ensure all of the

lingering oxygen in the chamber from the MgO growth has been pumped out. Once the heater

has cooled, the environment is adjusted for the Al growth which consists of 7 mTorr Ar with a

flow rate of 7 sccm. A working distance of 9 cm was used with a dc power of 50 W, which

resulted in sputter rate of 0.232 nm/s and a total deposition time of 7 minutes and 10 seconds.

49

The Al layer forms a polycrystalline layer and, as such, the RHEED pattern is very predictable

and ring like throughout the 100 nm deposition.

Following the trilayer deposition, the sample was removed from the chamber and ex situ

measurements were carried out. XRD measurements were performed to determine the

crystalline state of each layer and their epitaxial relationship. RHEED gives a basic look into the

crystal structure in real time but since it is only surface sensitive it, cannot probe the total

thickness of each layer which can evolve during the growth at high temperature. The out of

plane 2θ – θ scan can be seen in Figure 2.5a. From the wide range out of plane scan, all three

layers can be indexed in addition to the substrate which is (0001) oriented. Strong (0001) related

epitaxial peaks from the Re layer are observable as a result of Re’s high large atomic scattering

factor, 100 nm thickness, and good single crystalline quality. Additionally, peaks originating

from the MgO layer being (111) oriented as expected are seen and significantly weaker than the

Re layer given the smaller thickness and significantly lower z-number of both Mg and O leading

to a reduced atomic scattering factor. No other peaks can be indexed for MgO or Re, suggesting

that they are epitaxially grown and are well aligned to the substrate orientation. Multiple peaks

can be indexed to Al, confirming that the Al grows polycrystalline and not epitaxially.

The out-of-plane XRD confirms what was observed with RHEED in situ, but to further

investigate the in-plane crystallographic relationship between each layer an in-plane phi scan for

each layer is required. The relative orientations are described with respect to the substrate

orientation and, as such, first the in-plane (116) phi scan of the sapphire was measured as seen in

Figure 2.5 which, as expected, shows six peaks confirming the 6 fold symmetry of the hexagonal

sapphire. Following this, a phi scan of the (112) family of planes for Re was measured and once

again showed six peaks associated with the hexagonal crystal structure of Re. Interestingly the

50

peaks are shifted by 30° with respect to the substrate peaks indicating that the in-plane axes of

the Re layer is rotated 30° with respect to the substrate. When considering the (200) in-plane

Figure 2.5. Structural characterization. The out-of-plane XRD scan is shown in a with the film and

substrate peaks indexed. The Al polycrystalline peaks are indicated by squares and the sapphire substrate

peaks are shown with circles. The relative epitaxial relationship is demonstrated via phi scans in b-d for

the substrate, Re layer, and MgO layer respectively. A schematic of the orientation for each layer is given

in e-g where the twinned MgO film is depicted giving rise to the 6 in-plane phi peaks.

51

phi scan for MgO, one would expect to observe 3 peaks separated by 180° given the three-fold

symmetry that exists for (111) oriented cubic systems. However, in the (200) phi scan 6 peaks

60° apart were observed, which is consistent with the twinning structure that were predicted to

exist previously. Additionally, it can be seen that the peaks match the substrate orientation and

not the Re in-plane orientation. This confirms our previous model for the epitaxial relationship

of the three layers and their crystalline quality.

While the MgO based capacitor structures grew with fairly good crystalline quality they

did not perform well in terms of the energy loss. The lower Q value was determined to be

limited by the MgO layer and, as such, the possibility of using Al2O3 not only as the substrate

material but also the dielectric layer was explored. Other SSQCDs use Al2O3 as the dielectric

layer in the Josephson Junction via an amorphous deposition followed by a crystallizing

annealing step.19

However, this technique only works for very thin layers and is not ideal for the

shunted capacitor. For this work two separate growth processes were attempted: a two-step

process involving an amorphous layer followed by an anneal which served as a thin template

layer for the remainder of the growth and a one-step epitaxially deposition of Al2O3. The

remainder of the structure was identical to what was described previously with the bottom

electrode consisting of 100 nm of Re and the top electrode being a polycrystalline 100 nm Al

layer.

For the two-step process, having real time RHEED analysis was critical for being able to

identify when the amorphous layer had crystallized. The progression of the RHEED images

starting from the sapphire substrate can be seen in Figure 2.6. The previously seen sapphire

pattern shows a sharp RHEED pattern with Kikuchi lines, then following the Re growth the spots

are still defined but become slightly streaky and the Kikuchi lines disappear but are clearly still

52

single crystalline. Then 2 nm of Al2O3 were deposited by PLD at room temperature and the

associated RHEED pattern indicated that the film is amorphous due to the lack of features

besides the transmitted beam and Laue circle. This layer was deposited in 0.7 mTorr O2 with a

flow rate of 10 sccm. For 2 nm of growth, 400 pulses were needed. Following this thin layer, the

sample temperature was ramped to 850 °C in vacuum. Monitoring the RHEED pattern, the

diffraction streaks appear during the annealing process indicating a crystallization of the film

takes place and the surface is 2-dimensional. Once the crystallization occurs, the pressure is

returned to 0.7 mTorr O2 and the remaining 18 nm is deposited at high temperature on the seed

Figure 2.6. Two step Al2O3 RHEED progression is shown starting from the sapphire substrate

in a, followed by the Re layer in b. The first 2 nm of Al2O3 are grown at room temperature and

are completely amorphous as shown by the RHEED pattern in c. After the amorphous layer is

deposited, the temperature is raised to 850 °C where a crystallization is observed in d and the

subsequent RHEED pattern after 50 nm of Al2O3 is deposited in e indicating a smooth

crystalline surface.

53

layer to achieve a 20 nm total dielectric layer with the final RHEED pattern. This resulted in the

dielectric layer showing sharper spots indicating a higher degree of crystallinity.

The single step process follows a similar procedure compared with the two step except

the amorphous layer is skipped. Following the Re deposition, the temperature was held at 850

°C and after the Re post anneal the environment was changed to an argon one with 10 sccm

flowing Ar and a pressure of 0.8 mTorr. The Ar environment was chosen to avoid oxidizing the

Re at such a high temperature and rely solely on the Al2O3 target to provide enough oxygen for

the film. With these conditions, the RHEED pattern showed a spot like pattern throughout the

deposition.

Figure 2.7. Surface images of Al2O3 and Re. AFM image and RHEED pattern of the

grown Re and Al2O3 layers: (a) Epitaxial Re on Al2O3 substrate and (b) STM image of 100

nm-thick Re film showing single unit cell steps. The line scan of STM image of Re film. (c)

AFM image and RHEED pattern of epitaxial Al2O3 layer on Re/Al2O3 substrate.4

54

The AFM and STM images taken ex situ of the single step film can be seen in Figure

2.7,4 which show the Re surface after 100 nm deposition on the sapphire substrate and the

corresponding RHEED pattern. The STM image of the Re film shows the screw dislocations

that form and the hexagonal structure with one unit cell high steps. Additionally, the AFM

image from the Al2O3 is shown with its corresponding RHEED image. The Re and Al2O3 films

have an RMS roughness of 0.7 and 1.6 nm respectively with some indication of a granular island

structure.

For the single step processing films, TEM analysis was performed as a way to observe

the interface structure and film quality as seen in Figure 2.8.4 The full trilayer structure with

sharp interfaces can be seen in the low magnification image. A higher magnification image of

the Re/Al2O3 interface is also shown where no reaction layer exists even with the single step

processing and an atomically abrupt transmission exits. The planar selective area electron

diffraction (SAED) shows some elongation of the Al2O3 spots suggesting there is some mosaic

spread in the film with <7° dispersion consistent with some low angle grain boundaries. This

also confirms the epitaxial relationship in the XRD data.

Figure 2.8. TEM analysis of the heterostructure. (a) Low magnification cross-sectional TEM images

of Al/Al2O3/Re trilayer on (0001) Al2O3 substrate. (b) The cross-sectional HRTEM image near the

interface between Al2O3 thin film and Re layer. (c) Planar view selected area electron diffraction (SAED)

pattern of Al2O3 thin film and Re layer.4

55

Out-of-plane XRD scans were carried out on this multilayer stack and have similar

results to what was found in the MgO heterostructure with the Al polycrystalline peak and Re

(0002) peak clearly visible. However, for these samples the dielectric film peak from the

substrate peak cannot be distinguished as they are the same material and perfectly overlap. The

XRD scans can be seen in Figure 2.94 along with a Re rocking curve which measures the mosaic

spread and crystalline quality with a FWHM of 0.2°. The in-plane epitaxial relationship was

measured by XRD and matches what was previously shown for the Re layer and as such is not

duplicated here. As previously discussed, the dielectric layer matches that of the substrate and

Figure 2.9. XRD of the Al/Al2O3/Re trilayer. (a) X-ray diffraction θ-2θ scan of the Al/Al2O3/Re

trilayer on (0001) Al2O3 substrate. (b) Rocking curve θ-scan of the (0002) Re reflection. (c) X-ray

azimuthal ϕ-scans of the Al2O3 (11�̅�6) and Re (11�̅�4) reflections. (d) Schematic diagram of [2𝟏𝟏̅̅̅̅ 0]

Al2O3//[10�̅�0] Re in-plane relationship between Al2O3 substrate and Re film.4

56

the phi scan does not show any new discernable peaks indicating the Al2O3 layer is aligned in-

plane to the substrate.

Once high quality multilayer capacitive structures were achieved consisting of epitaxial

bottom electrodes and dielectric layers, they were further characterized by measuring their

internal Q value as a function of voltage. The processing results and their impact, while this is

outside the scope of this thesis, are an important way to determine the performance of these

heterostructures and will be briefly discussed here. Two separate measurements were

performed: one to measure the performance of the full Al/Al2O3/Re/sapphire heterostructure and

the other to only measure the Q from only the Re deposited on sapphire. In this way the loss

associated with the dielectric layer can be decoupled from the Re layer. For these measurements,

two different types of resonators were patterned on the samples: a lumped element LC resonator

Figure 2.10. Quality factor measurements of the

trilayer. (a) Layer stack of the lumped element

LC resonator device showing epitaxial trilayer and

SiN x overlap coupling capacitor connected to the

measurement feedline. The thicknesses for trilayer

stack were 100 nm, 20 nm, and 100 nm for Re,

Al2O3, and Al, respectively. (b) Resonator CAD

layout. (c) Electrical circuit schematic of the

resonator. (d) Internal loss 1/Q i of epitaxial

Re/Al2O3/Al LC resonators and CPW Re

resonators versus rms voltage across the resonator.

The LC resonators incorporate trilayers grown

according to both the one-step and two-step

processes described in the text; the multiple

datasets for the one-step growth and for the CPW

Re represent different growth and fabrication runs

with nominally identical parameters.4

57

for the full multilayer and a coplanar waveguide (CPW) resonator for the Re layer. The final

material stack and resonator structure can be seen in Figure 2.10.4 Once patterned, the samples

were cooled to 50 mK using an adiabatic demagnetization refrigerator. The two-step and one-

step grown Al2O3 dielectric layer in this way can be directly compared. The two-step multilayer

shows an internal Q of 2 x 104 while the Q measured for the one step is 3 x 10

4 indicating that

the one-step processing yields a higher quality dielectric layer and shows comparable intrinsic Q

to the best amorphous Si:H dielectrics.24

The CPW resonator measurements of the Re layer

shows similar values to the full heterostructure which suggests that the Re layer Q value may be

the limiting material for the shunt capacitor and that the internal Q of the Al2O3 layer may in fact

be higher than measured here.

2.4. Conclusions

Further improvement in Re electrode layer quality is required to advance the progress of

these multilayer stacks. Nonetheless, these measurements confirm what was observed in situ

with RHEED. High crystalline quality films are formed and interfaces preserved and single

crystalline Al2O3 can be grown at high temperature without degrading the Re layer by oxidizing

it. The benefit of having both sputtering and PLD in one system coupled with nearly continuous

RHEED monitoring throughout the entire trilayer deposition is vast. The information that was

extracted about interface quality and the crystalline state in real time saved countless hours of

wasteful film deposition and ex-situ analysis, clearly demonstrating the power of the Hybrid

system.

58

2.5. References for chapter 2

1 Niranjan, M. K., Duan, C.-G., Jaswal, S. S. & Tsymbal, E. Y. Electric field effect on

magnetization at the Fe/MgO(001) interface. Applied Physics Letters 96, 222504,

doi:10.1063/1.3443658 (2010).

2 Heron, J. T. et al. Deterministic switching of ferromagnetism at room temperature using

an electric field. Nature 516, 370-373, doi:10.1038/nature14004 (2014).

3 Makhlin, Y., Schon, G. & Shnirman, A. Josephson-junction qubits with controlled

couplings. Nature 398, 305-307, doi:10.1038/18613 (1999).

4 Cho, K. H. et al. Epitaxial Al2O3 capacitors for low microwave loss superconducting

quantum circuits. Apl Mater 1, 042115, doi:Artn 042115 10.1063/1.4822436 (2013).

5 Martinis, J. M. et al. Decoherence in Josephson qubits from dielectric loss. Phys Rev Lett

95, 210503, doi:10.1103/PhysRevLett.95.210503 (2005).

6 Schoelkopf, R. J. & Girvin, S. M. Wiring up quantum systems. Nature 451, 664-669,

doi:10.1038/451664a (2008).

7 Phillips, W. A. Two-level states in glasses. Reports on Progress in Physics 50, 1657

(1987).

8 Gao, J. S., Zmuidzinas, J., Mazin, B. A., LeDuc, H. G. & Day, P. K. Noise properties of

superconducting coplanar waveguide microwave resonators. Applied Physics Letters 90,

102507, doi:10.1063/1.2711770 (2007).

9 Gao, J. S. et al. Experimental evidence for a surface distribution of two-level systems in

superconducting lithographed microwave resonators. Applied Physics Letters 92, 152505,

doi:10.1063/1.2906373 (2008).

59

10 Alford, N. M. et al. Dielectric loss of oxide single crystals and polycrystalline analogues

from 10 to 320 K. Journal of the European Ceramic Society 21, 2605-2611,

doi:10.1016/s0955-2219(01)00324-7 (2001).

11 Chrisey, D. B. & Hubler, G. K. Pulsed Laser Deposition Of Thin Films. (John Wiley &

Sons, Inc, 1994).

12 Christen, H. M. & Eres, G. Recent advances in pulsed-laser deposition of complex

oxides. Journal Of Physics-Condensed Matter 20, 264005, doi:10.1088/0953-

8984/20/26/264005 (2008).

13 Doeswijk, L. M., Rijnders, G. & Blank, D. H. A. Pulsed laser deposition: metal versus

oxide ablation. Appl. Phys. A-Mater. Sci. Process. 78, 263-268, doi:10.1007/s00339-003-

2332-0 (2004).

14 Abraham, M. M. Growth of High-Purity and Doped Alkaline Earth Oxides: I. MgO and

CaO. The Journal of Chemical Physics 55, 3752, doi:10.1063/1.1676658 (1971).

15 Maksym, P. A. & Beeby, J. L. A Theory of Rheed. Surface Science 110, 423-438,

doi:10.1016/0039-6028(81)90649-X (1981).

16 Rijnders, G. J. H. M., Koster, G., Blank, D. H. A. & Rogalla, H. In situ monitoring during

pulsed laser deposition of complex oxides using reflection high energy electron

diffraction under high oxygen pressure. Applied Physics Letters 70, 1888-1890,

doi:10.1063/1.118687 (1997).

17 Svedberg, E. B., Birch, J., Edvardsson, C. N. L. & Sundgren, J. E. Real time

measurements of surface growth evolution in magnetron sputtered single crystal Mo/V

superlattices using in situ reflection high energy electron diffraction analysis. Surface

Science 431, 16-25, doi:10.1016/S0039-6028(99)00498-7 (1999).

60

18 Herklotz, A. et al. Stoichiometry control of complex oxides by sequential pulsed-laser

deposition from binary-oxide targets. Applied Physics Letters 106, 131601,

doi:10.1063/1.4916948 (2015).

19 Oh, S. et al. Elimination of two level fluctuators in superconducting quantum bits by an

epitaxial tunnel barrier. Physical Review B 74, 100502, doi:10.1103/Physrevb.74.100502

(2006).

20 Oh, S. et al. Epitaxial growth of rhenium with sputtering. Thin Solid Films 496, 389-394,

doi:10.1016/j.tsf.2005.09.091 (2006).

21 Wang, H. et al. Improving the coherence time of superconducting coplanar resonators.

Applied Physics Letters 95, 233508, doi:10.1063/1.3273372 (2009).

22 Tanuma, S., Powell, C. J. & Penn, D. R. Calculations of electron inelastic mean free

paths. IX. Data for 41 elemental solids over the 50 eV to 30 keV range. Surface and

Interface Analysis 43, 689-713, doi:10.1002/sia.3522 (2011).

23 Ichimiya, A. & Cohen, P. I. Reflection High-Energy Electron Diffraction. (Cambridge

University Press, 2004).

24 O’Connell, A. D. et al. Microwave dielectric loss at single photon energies and

millikelvin temperatures. Applied Physics Letters 92, 112903, doi:10.1063/1.2898887

(2008).

61

3. Two-dimensional electron gas (2DEG) at the LaAlO3/SrTiO3

interface by sputtering

3.1. Introduction

In this chapter we will discuss one example of the use of in situ RHEED with sputtering.

Oxide heterostructure have the vast potential to transform electronic devices due to their high

degree of tunability, functionality, and stability in extreme environments.1 One oxide

heterostructure, LaAlO3 (LAO) deposited on SrTiO3 (STO), has garnered a lot of attention since

2004 when it was discovered that there is a conducting state at the interface between the two

band insulators, known as a two-dimensional electron gas (2DEG).2,3

In order to obtain this

conducting interface, the thickness of the LAO must be very precisely controlled, leading to

extremely sensitive growth conditions; therefore, the vast majority of LAO/STO systems are

grown in systems that have incorporated in situ RHEED, most commonly pulse laser deposition

(PLD) which is typically not technologically scalable. Our work,4 which demonstrates for the

first time a conducting interface at the LAO/STO interface grown by sputtering, was critical for

the LAO/STO system to become commercially viable as well as all other equally sensitive oxide

heterostructures. This chapter details the background of LAO/STO and how the incorporation of

in situ RHEED with sputtering has improved its potential of being used on an industrial scale.

Both LAO and STO have the perovskite ABO3 structure with similar enough lattice

parameters that LAO can be grown epitaxially on STO. When LAO is grown specifically on

62

TiO2 terminated STO substrates5 in low partial pressures of oxygen (~10

-6 torr) a metallic like

state forms.2 In this work by Ohtomo and Hwang in 2004 they demonstrate that without the

TiO2 termination a conducting interface does not form. Subsequently when the conducting

interface, which is often referred to as a two-dimensional electron gas or 2DEG, is formed they

show that as the partial pressure of oxygen in the growth chamber increases the conductance and

mobility of the 2DEG degrades.

Nakagawa et al proposed the polar catastrophe model for this system as shown in Figure

3.1.6 The main mechanism that governs this system is a polar discontinuity that occurs when

you take a material such as SrTiO3 which has alternating layers of SrO and TiO2 both of which

have a net charge of zero in the (001) orientation. The discontinuity begins when LaAlO3 is

deposited epitaxially on top of the STO. The LAO has repeating layers of LaO and AlO2 which

have net charges of +1 and -1 respectively. The polar nature of the LAO on top of the non-polar

STO creates a polar discontinuity. The result of this is a diverging potential as a function of

LAO thickness. As seen in Figure 3.1a-b, the termination of the STO substrate directly affects

the layering order on top of it, which in turn changes the sign of the diverging potential from

positive in a) to negative in b). In this system, the diverging potential will quickly become

unstable. To avoid this diverging potential the system distributes either half an electron to the

interface or to the surface depending on the substrate termination. This mitigates the diverging

potential and causes it to oscillate around zero and is known as the “polar catastrophe model”.

When the substrate is TiO2 terminated, half an electron moves to the interface which provides the

carriers required for a conducting interface to form.

63

As discussed previously, a diverging potential as a function of thickness forms due to the

polar discontinuity. In 2006, Thiel et al.7 showed that the LAO/STO system can maintain the

diverging potential up until 4 unit cells at which point the polar catastrophe model takes effect.

This is clearly seen in Figure 3.2 where no charge carriers are present in the heterostructure up

until 4 unit cells at which point the carrier concentration rapidly jumps up and then is relatively

constant as a function of number of unit cells. This critical thickness of 4 u.c. has become a

universal standard for all 2DEGs formed at the (001) LAO/STO interface and can be viewed as a

telltale sign that a 2DEG has formed and the conduction is not simply coming from oxygen

vacancies in the STO or other defects.

The observation of RHEED oscillations in MBE growth of classical semiconductors has

been around for many years8,9

and then later introduced to oxide MBE growth as well.10

High-

pressure RHEED was initially introduced as an in situ analysis technique for PLD by Rijnders et

al11

in 1997 and since then has proven to be an essential tool for the growth of epitaxial oxide

Figure 3.1. Layering structure in the LAO/STO heterointerface with TiO2 terminated STO interface

in a and SrO termination in b and their associated diverging potentials. C and d show how the

redistribution of charges to the interface can mitigate the diverging potentials to form a stable state.6

64

films and control over complex interfaces. The LAO/STO system is a perfect example of this

due to the critical thickness effect; having precise control of the thickness has a significant

impact on the electronic properties of the interface.

Despite the prominent position of a 2DEG formation at the LAO/STO interface in the

oxide research community since its discovery in 2004, there is still a large amount of discussion

surrounding the origin of the carriers at the interface. In addition new properties are

continuously emerging with respect to this material system. In order to expand this field to a

broader community, a more readily available deposition technique such as sputtering should be

pursued. Using sputtering as a new growth technique for depositing epitaxial LAO on STO

substrates can create a new avenue for scientific studies and bring further insights to how the

2DEG is formed.

Figure 3.2. Carrier concentration as a function of number of unit

cells where a clear increase in carriers is observed at 4 unit cells and

below which no carriers are detectable.7

65

As previously mentioned, work has already been done with a focus on creating devices

using the LAO/STO system12

and in addition to this it has been shown that the 2DEG can be

formed when both STO and LAO are grown on Si substrates13

demonstrating that the LAO/STO

interface can be incorporated into current industrial platforms and has technological relevance.

However, in order for these devices to be broadly applicable they need to be integrated with

current fabrication technology—most notably there is a need to deposit these thin films

uniformly on large wafers. 90° off-axis sputtering is a scalable process that has been shown to

create smooth, epitaxial films that are uniform over a large area.14,15

This is essential for

integration with silicon. While several groups have grown LAO/STO heterostructures by sputter

deposition,16,17

none have reported a conducting interfacial 2DEG. This task requires finely

tuned growth parameters that avoid having the bulk substrate conducting and the entire sample

insulating, while maintaining a 2DEG at the interface.

Here, we report on the steps necessary for the creation of a 2DEG at the LaAlO3/SrTiO3

heterointerfaces grown by 90° off-axis sputtering.4 The objective is to show the electrical

transport properties of the LaAlO3/SrTiO3 heterointerfaces grown by sputtering are comparable

to those grown by pulsed laser deposition. This would indicate that sputtering can be a viable

alternative to the typical PLD growth of these interfaces. Demonstrating room-temperature

conductive-atomic force microscope switching of 2DEG nanostructures would further show that

functional applications of this material system are extended to the scalable growth technique of

sputtering.

66

3.2. Film growth and structural characterization

In order to create an oxide interfacial 2DEG, we use a different regime of growth

conditions, deviating from the previously reported 3:4 O2:Ar ratio16,18

and high pressure O2

sputter environments.17

In the case of PLD, it has been shown that too high (greater than 7.5

mTorr),19

or too low (less than 7.5 x 10-7

mTorr)20

oxygen pressure can degrade LAO structural

quality and inhibit 2DEG formation. Even though oxygen vacancies are known to be removed

during post-annealing after LAO growth, earlier works have shown that low oxygen partial

pressure can play a major role in the creation of a 2DEG.21,6,22

However, duplicating the 10-4

–

10-6

Torr21

O2 partial pressures used in PLD can be challenging in sputtering due to typically

higher Ar gas pressure (i.e. 200 mTorr) used for 90° off-axis sputtering. There is a 3-5 order of

magnitude difference in partial pressures between Ar and O2, and pressure in this regime cannot

easily be controlled accurately.

We overcome this difficulty by sputtering in pure Ar and relying on the sputtered single

crystal LAO target to sustain a large enough background pressure of O2 and atomic oxygen to

create an environment similar to PLD. During sputtering of an oxide target material, a relatively

large amount of O2 and more importantly atomic oxygen is produced. The nearly zero activation

energy associated with atomic oxygen compared with the non-zero energy of molecular oxygen23

significantly increases the effective PO2 during sputter growth compared with PLD where only

molecular oxygen is present. This mechanism creates an environment that can produce high

quality 2DEGs2,24

during sputtering while only introducing Ar gas.

Our growth system consists of a 2 inch RF magnetron sputter source in a 90° off-axis

geometry with reference to the sample heater. All films discussed here were grown on TiO2-

terminated (001) STO substrates.5 A LAO/STO control sample was grown by PLD using

67

conditions reported elsewhere.13

Currently pulsed laser deposition (PLD)2,21

is the dominant

deposition technique used for creating oxide interfacial 2DEG heterostructures, and only more

recently has molecular beam epitaxy (MBE)25,26

been used to create a 2DEG. Both of these

growth techniques take advantage of reflection high-energy electron diffraction (RHEED) as an

in situ diagnostic tool, enabling layer by layer control at the unit cell level.8-11

As such, a PLD

grown sample was used as a control. The sputtered samples were grown from a 2-inch single

crystal LAO target mounted on a US Gun II sputter source at an RF power of 50 W. The sample

temperature was 780 °C during growth, consistent with previously reported works.21

A partial

pressure 200 mTorr of Ar was used, with a minimum pre-sputter time of 15 minutes in order to

stabilize a background partial pressure of O2 (and atomic oxygen) produced from the target. This

is a vital step in the growth process as films grown without this extended pre-sputtering were

found to be insulating. It should also be noted that samples we grow at higher PO2, for example

3:4 of O2:Ar, produce insulating films as well. We subsequently annealed the sample for 1 hour

in 300 Torr O2 at 600 °C to ensure that the bulk substrate was not conducting.27

We have utilized x-ray reflectivity as a thickness calibration; this is essential since the

thickness of LAO directly modifies the electronic characteristics of the LAO/STO interface.7 At

the deposition conditions described above, we are able to grow LAO films at a rate of 1.65

Å/min as determined by fitting x-ray reflectivity data as shown in Figure 3.3a. This data also

indicates that the film is smooth with small roughness.

68

Further x-ray analysis shown in Figures 3.3b-d confirms the high-quality epitaxial growth

of the LAO thin films. These high-resolution x-ray diffraction results from a 40 nm sample

(Figures 3.3b-c) show an out-of-plane lattice parameter of 3.74 Å (Figure 3.3b) and in plane

lattice parameter of 3.89Å. This is close to the 3.905 Å STO in-plane lattice constant of the

substrate, leading to a 2.3% tensile strain in the film, referenced to a bulk LAO lattice constant of

3.790Å. This indicates that even at 40nm the film is almost fully strained. The azimuthal phi

scan in Figure 3.3c shows 4-fold symmetry of the LAO film with an in-plane (101) FWHM of

0.26° indicating a very small in-plane mosaic spread of the film while the in-plane FWHM of the

Figure 3.3. Structural characterization of the LAO/STO heterostructure. a) x-ray

reflectivity measurement of the LAO/STO heterostructure displaying clear thickness

plateaus closely matching the model. b) Wide range two theta-omega scan showing

single crystal LAO grown on a STO substrate. c) In-plane (101) phi scan showing 4-

fold symmetry. d) Reciprocal space map of a 10 unit cell sample showing the LAO film

to be fully coherent.

69

substrate is 0.01°. As shown in Figure 3.3d, a reciprocal space mapping (RSM) around the (1̅03)

Bragg peak of a STO substrate shows that the LAO film is fully coherent and of single phase.

The elongation of the (1̅0L) film peak arises from the ultra-thin thickness of only 10 unit cells

LAO. The out-of-plane lattice constant obtained from the peak position of the LAO film is 3.72

Å, which is in good agreement with that measured from the θ-2θ scan on the thicker film.

The sputtered LAO film is atomically smooth with an average roughness of ~0.14 nm,

as shown in the atomic force microscopy (AFM) images of Figure 3.4. The clear step and terrace

structures observed in Figure 3.4a from the treated STO substrate persist in the LAO film on top

as seen in Figure 3.4b.

Figure 3.4. Surface topography of the STO substrate and LAO film in a and b

respectively. The AFM shows the 10 unit cell film following the smooth step and

terrace structure of the substrate. A plot of height versus distance along the line

showing clear unit cell high steps is shown under each AFM image.

70

3.3. Electric characterization

A study of the carrier concentration (ns) as a function of thickness was carried out to show

the existence of a critical thickness. Samples were wire bonded with Al wires in the four-point

van der Pauw geometry and measured Hall coefficients were found to be linear as a function of

applied magnetic field. Figure 3.5a shows a clear transition between a conducting state and an

insulating state at 4 unit cells, consistent with findings on PLD-grown films.7 There can also be

seen a negligible thickness dependence for thicknesses greater than 4 unit cells. This is a clear

Figure 3.5. Transport properties of the sputtered LAO/STO heterostructure. a) The carrier

concentration (ns) as a function of LAO film thickness. Plot of b) carrier concentration (ns), c) mobility

(μH), and d) sheet resistance (RS) as a function of temperature of a 10 unit cell sputter sample compared

with a PLD sample. Blue dashed lines in b) represent power law fits to the mobility above 75 K.

71

indication that these LAO/STO samples produced by sputtering exhibit a 2DEG behavior as

explained by electronic reconstruction.

Additionally, transport measurements were carried out as a function of temperature on both

sputtering and PLD samples with similar growth conditions for comparison with LAO

thicknesses of 10 unit cells in the sputtered sample. The carrier density in the sputtered sample

was found to have a somewhat weaker dependence on temperature than the PLD grown sample

(Figure 3.5b). The temperature dependence of the Hall mobility (μH) is similar above 75 K,

where both obey a power law behavior of T-α

, with α = 2.3 and 2.4 for the sputtered and PLD

grown sample respectively (Figure 3.5c), consistent with reported behaviors.20,24

At low

temperatures the mobility in the sputtered sample saturates at lower values compared with the

PLD grown sample, suggesting increased scattering at these growth conditions. The increased

low-temperature sheet resistance (Rs) is consistent with the reduced mobility values since the

variation in carrier density is small in this range; in the range above 50 K both samples display

similar metallic behavior (Figure 3.5d).

We also demonstrate room-temperature conductive-AFM (c-AFM) switching of 2DEG

nanostructures formed at LAO/STO heterointerfaces grown by 90° off-axis sputtering. The

ability to write and erase nanostructures in these samples has been observed at a range of

thicknesses between 3 and 3.7 unit cells. After 90° off-axis sputtering growth, electrically

conducting contacts to the interface are defined by optical lithography. Then the sample is

prepared by initially milling 25 nm deep trenches through the LAO layer via an Ar-ion mill and

then filling them with 4 nm of Ti followed by 25 nm of Au to form bilayer electrodes via

sputtering. Within the 30 μm × 30 μm “canvas” defined by the electrical contacts, nanostructures

are written and erased at the interface using c-AFM lithography at room temperature. The

72

conductance between two electrical contacts is continuously monitored by a lock-in amplifier

(Figure 3.6a, b). Figure 3.6 shows a typical c-AFM writing and erasing process in a 3.4 unit cell

LAO/STO sample with an initially insulating interface. To begin with, two rectangular pads are

“written” (that is, raster scanning by a Vwrite= +10 V biased AFM tip) for better contacting to the

two Au electrodes, then a conductive nanowire is created by the +10 V biased AFM tip scanning

from one electrode to another at 300 nm/s speed (Figure 3.6a). When the tip reaches the other

electrode, a pronounced and abrupt conductance jump is observed (Figure 3.6c). The observed

Figure 3.6. Room-temperature conductive-AFM (c-AFM) switching of sputtered LAO/STO.

a) Schematic diagram of the writing process. AFM tip is +10V biased. b) Schematic diagram of

the erasing process. AFM tip is -10V biased and moving perpendicularly across the nanowire. c)

Conductance between the two electrodes during the writing process. When the tip reaches the

second electrode, conductance increases abruptly. d) Conductance between the two electrodes

during the erasing process. As the tip scans across the nanowire, conductance decreases to zero.

The red curve shows the best fit indicating a nanowire thickness of 9.2 nm.

73

maximum conductance change and non-exponential decay in atmosphere conditions are

comparable with the 3.4 unit cell LAO/STO samples grown by PLD.12

After writing the nanowire, the AFM tip is repositioned and biased at Verase= -10 V, then

moving perpendicularly across the nanowire at 10 nm/s speed (Figure 3.6b). The conductance

decreased abruptly to zero when the tip reaches the nanowire (Figure 3.6d). The nanowire width

can be qualified by fitting the conductance drop curve with the function G(x) = G0-G1tanh(x/h),12

the red curve in Figure 3.6d. From the best fitting parameters we get a nanowire width of 9.2

nm. The ability to create conductive nanostructures is important for technological applications,

but it also serves as a sensitive probe of the uniformity of the 2DEG. A single insulating patch

along the nanowire is sufficient to prevent conductivity. Here, the properties are comparable to

PLD-grown heterostructures.12

3.4. Conclusions

Sputtering is a thin film growth technique capable of the uniform large area deposition

required for scaling. In this work we have demonstrated that sputtering can be used to grow high

quality epitaxial LaAlO3 films with excellent surface quality on TiO2-terminated SrTiO3

substrates. Growth conditions were chosen in attempt to mimic those used by growth processes

known to form a 2DEG at the interface, and this resulted in the creation of an LaAlO3/SrTiO3

two-dimensional electron gas made by sputtering. We showed that these samples have

comparable electronic transport properties to heterostructures grown with pulsed laser

deposition. We have also demonstrated room-temperature conductive-AFM switching of these

2DEG nanostructures showing the capability to make interesting oxide devices on these films.

74

In summary, we have created another avenue of exploration for the study of the 2DEG at the

LaAlO3/SrTiO3 interface by developing a growth process using a sputter deposition method.

3.5. References

1 Mannhart, J. & Schlom, D. G. Oxide Interfaces—An Opportunity for Electronics.

Science 327, 1607-1611, doi:10.1126/science.1181862 (2010).

2 Ohtomo, A. & Hwang, H. Y. A high-mobility electron gas at the LaAlO3/SrTiO3

heterointerface. Nature 427, 423-426, doi:10.1038/nature02308 (2004).

3 Ohtomo, A., Muller, D. A., Grazul, J. L. & Hwang, H. Y. Artificial charge-modulationin

atomic-scale perovskite titanate superlattices. Nature 419, 378-380,

doi:10.1038/nature00977 (2002).

4 Podkaminer, J. P. et al. Creation of a two-dimensional electron gas and conductivity

switching of nanowires at the LaAlO3/SrTiO3 interface grown by 90° off-axis sputtering.

Applied Physics Letters 103, 071604, doi:10.1063/1.4817921 (2013).

5 Kawasaki, M. et al. Atomic Control of the SrTiO3 Crystal Surface. Science 266, 1540-

1542, doi:10.1126/science.266.5190.1540 (1994).

6 Nakagawa, N., Hwang, H. Y. & Muller, D. A. Why some interfaces cannot be sharp. Nat

Mater 5, 204-209, doi:10.1038/nmat1569 (2006).

7 Thiel, S., Hammerl, G., Schmehl, A., Schneider, C. W. & Mannhart, J. Tunable quasi-

two-dimensional electron gases in oxide heterostructures. Science 313, 1942-1945,

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8 Wood, C. E. C. & Joyce, B. A. Tin‐doping effects in GaAs films grown by molecular

beam epitaxy. Journal of Applied Physics 49, 4854-4861, doi:10.1063/1.325517 (1978).

9 Harris, J. J., Joyce, B. A. & Dobson, P. J. Oscillations in the Surface-Structure of Sn-

Doped Gaas during Growth by Mbe. Surface Science 103, L90-L96, doi:10.1016/0039-

6028(81)90091-1 (1981).

10 Bozovic, I. & Eckstein, J. N. Analysis of Growing Films of Complex Oxides by Rheed.

MRS Bulletin 20, 32-38, doi:10.1557/S0883769400044870 (1995).

11 Rijnders, G. J. H. M., Koster, G., Blank, D. H. A. & Rogalla, H. In situ monitoring during

pulsed laser deposition of complex oxides using reflection high energy electron

diffraction under high oxygen pressure. Applied Physics Letters 70, 1888-1890,

doi:10.1063/1.118687 (1997).

12 Cen, C., Thiel, S., Mannhart, J. & Levy, J. Oxide Nanoelectronics on Demand. Science

323, 1026-1030, doi:10.1126/science.1168294 (2009).

13 Park, J. W. et al. Creation of a two-dimensional electron gas at an oxide interface on

silicon. Nat Commun 1, 94, doi:10.1038/ncomms1096 (2010).

14 Eom, C. B. et al. In situ grown YBa2Cu3O7- thin films from single-target magnetron

sputtering. Applied Physics Letters 55, 595-597 (1989).

15 Eom, C. B., Marshall, A. F., Laderman, S. S., Jacowitz, R. D. & Geballe, T. H. Epitaxial

and Smooth Films of a-Axis YBa2Cu3O7. Science 249, 1549-1552,

doi:10.1126/science.249.4976.1549 (1990).

16 Lee, A. E. et al. Epitaxially grown sputtered LaAlO3 films. Applied Physics Letters 57,

2019-2021 (1990).

76

17 Dildar, I. M. et al. Conductivity of LaAlO3/SrTiO3 Interfaces made by Sputter

Deposition. CORD Conference Proceedings (2011).

18 Sader, E., Schmidt, H., Hradil, K. & Wersing, W. Rf-Magnetron Sputtered Lanthanum

Aluminate Buffer Layers on Silicon. Supercond Sci Tech 4, 371-373, doi:10.1088/0953-

2048/4/8/010 (1991).

19 Maurice, J.-L. et al. Electron energy loss spectroscopy determination of Ti oxidation state

at the (001) LaAIO3 /SrTiO3 interface as a function of LaAIO3 growth conditions. EPL

(Europhysics Letters) 82, 17003 (2008).

20 Kalabukhov, A. et al. Effect of oxygen vacancies in the SrTiO3 substrate on the electrical

properties of the LaAlO3 / SrTiO3 interface. Physical Review B 75, 121404 (2007).

21 Huijben, M. et al. Structure-Property Relation of SrTiO3/LaAlO3Interfaces. Advanced

Materials 21, 1665-1677, doi:10.1002/adma.200801448 (2009).

22 Reyren, N. et al. Superconducting interfaces between insulating oxides. Science 317,

1196-1199, doi:10.1126/science.1146006 (2007).

23 Yamamoto, K. et al. Role of Atomic Oxygen Produced by an Electron-Cyclotron

Resonance Plasma in the Oxidation of Yba2Cu3O7-X Thin-Films Studied by Insitu

Resistivity Measurement. Applied Physics Letters 57, 1936-1938, doi:10.1063/1.104147

(1990).

24 Huijben, M. et al. Structure-Property Relation of SrTiO3/LaAlO3 Interfaces. Advanced

Materials 21, 1665-1677, doi:10.1002/adma.200801448 (2009).

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in two-dimensional superconductors. Nat Phys 7, 895-900, doi:10.1038/nphys2075

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LaAlO3/SrTiO3 interfaces. Nat Commun 4, 2351, doi:10.1038/ncomms3351 (2013).

27 Moos, R., Menesklou, W. & Hardtl, K. H. Hall-Mobility of Undoped N-Type Conducting

Strontium-Titanate Single-Crystals between 19-K and 1373-K. Appl. Phys. A-Mater. Sci.

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78

4. In situ RHEED during oxide sputtering

4.1. Introduction

As the oxide field progresses, atomic layer controlled growth is becoming ever more

essential for the understanding and engineering of these complex oxide heterointerfaces. This

includes the understanding and manipulation of surfaces and interfaces. Currently, molecular

beam epitaxy (MBE) and pulsed laser deposition (PLD) techniques take advantage of reflection

high-energy electron diffraction (RHEED) as an in situ diagnostic tool of the structure of the

surface during deposition, enabling layer-by-layer control at the unit cell and sub unit cell level.1-

5 The observation of intensity oscillations of the RHEED specular reflection in MBE growth of

semiconductors has been exploited for many years to control stoichiometry and growth rate.1-3

The primary advantage of incorporating RHEED analysis to the growth of artificially layered

superlattices is the precise thickness control of each layer.6 The RHEED technique was readily

adapted to growth of complex oxides with MBE4 because the pressure during growth is

sufficiently low (<10–6

Torr) to avoid scattering of the electron beam and damage to the filament

in the electron source. Subsequently, RHEED at high oxygen pressures (<0.3 Torr) was

developed for PLD by Rijnders et al5 and has been widely adopted for growing epitaxial oxide

films and controlling complex interfaces by this technique.

The development of in situ RHEED analysis for sputter growth would introduce similar

advantages – for example, rapid optimization of growth parameters and control of growth rates,

79

and enhance reproducibility of interface and superlattice growth – to this widely used and

technologically important deposition technique.7-11

Moreover, analysis of in situ RHEED

intensities can provide fundamental information on epitaxial growth mechanisms (e.g. Stranski-

Krastanov, Frank-van der Merwe or Volmer-Weber modes12,13

and layer-by-layer versus step

flow) that is currently unknown for many thin film systems deposited by the sputter technique.

The concept behind layer-by-layer control by RHEED oscillations can be seen clearly in

Figure 1.6. As discussed in Chapter 1, the intensity of the RHEED specular spot is in direct

correlation to the percent of total coverage on a layer.14

When there is complete coverage the

intensity is at a maximum, and when there is only fifty percent coverage the intensity is at a

minimum due to scattering of the electron beam from a rough surface. Following this idea,

RHEED intensity versus time can be used to precisely control the number of unit cells deposited

during layer by layer growth. Beyond this, RHEED provides real time information regarding the

crystallinity, surface roughness as well as many other surface features.

This type of control becomes especially important in the LAO/STO system when

considering the ability to write and erase nanowires on the samples as first reported by Cen et al

in 2009.15

Here, they reported on using a conductive Atomic Force Microscopy (AFM) tip to

effectively “draw” conducting nanowires into the LAO/STO interface. This is only possible

when your films are very thin and below the critical thickness, then by applying a bias, an

additional field is generated in order to induce carrier movement to the interface causing the

interface to locally become conducting. We demonstrated this read and write ability in the

previous section with our sputter grown films but would have benefitted from the use of RHEED

during sputtering.

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4.2. RHEED with magnetron sputtering

Sputter deposition is a widely used growth technique for a large range of technologically

important material systems. Phase-pure and smooth epitaxial films of carbides,16

nitrides,17,18

oxides,7,19,20

metals,21,22

and more can be deposited by sputtering, from the research level up to

an industrial scale. As increased attention is given to design of emergent phenomena at

interfaces between these materials systems, atomic-layer controlled growth of epitaxial thin films

and artificially layered superlattices has become critical. Despite the many advantages of sputter

deposition, it suffers from a lack of available in situ analysis techniques for atomic-layer control

during growth. Reflection high-energy electron diffraction (RHEED) is generally used to

monitor thin film deposition yet it has not been commonly used during sputtering. This is largely

due to large magnetic fields from the sputter sources that strongly deviate the RHEED beam

trajectory. Additionally, typically high pressures of oxygen and argon associated with sputter

deposition of epitaxial complex oxides (>100 mTorr) can cause the electron beam to scatter

creating a large diffuse background of the intensity and reducing the overall detection range.

The goal of this work is to expand the functionality of this important in situ analysis

technique beyond MBE and PLD by applying it to sputter deposition. In this thesis we solve the

magnetic field problem by presenting a novel approach to minimize or eliminate the influence of

magnet fields in a wide range of growth geometries. The prediction and modeling of the

magnetic field effect on the electron beam is the focus of the following chapter while this chapter

will be more dedicated to the observed magnetic field effect, the gas effect, and mostly on the

use of RHEED in real-time during sputtering.

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In this chapter we demonstrate digital control of magnetron sputter deposition using in situ

high-pressure RHEED by applying this technique to the widely studied model oxide system,

SrRuO3 (SRO). During 90° off-axis sputtering of SRO films we observed strong specular spot

oscillations. This allows us to identify the growth mode as layer by layer and establishes our

ability to have unit cell control during sputter growth for the first time. Similar results are

observed and presented during the growth of perovskites La0.7Sr0.3MnO3 (LSMO) and LaAlO3

(LAO), confirming that this approach can be universally applied to sputter deposition of other

materials.

4.2.1. Scattering due to gas

As previously mentioned, this use of RHEED with sputtering is very challenging due to

the strong magnetic fields and the high partial pressures of Ar and other process gases present

during the deposition. Both of these have undesirable impacts on the electron beam used for

RHEED. The presence of process gas can scatter a significant portion of the electron beam.

This effect increases the overall background signal observed on the RHEED signal which

decreases the overall dynamic range that can be measured. In this section we aim to demonstrate

the effects of the variable gas environments and demonstrate how the sputter gas environment

differs from that of a typical oxide PLD environment and how that difference affects the RHEED

pattern. Following this we will provide some insight into ways to limit the effects from the

process gas and improve the measureable RHEED pattern.

Both PLD and sputtering have been reported to have working pressures that cover a huge

range of gas pressures, in particular PLD can go from UHV to hundreds of mTorr whereas the

82

lower limit for sputtering is several mTorr due to the need to have atomic collisions in order to

form a plasma. For 90° off-axis sputtering, gas pressures are typically in the hundreds of mTorr

range in order to reduce the number of high energy particles that reach the substrate by

shortening the mean free path and increasing the number of collisions. Sputtering

characteristically uses an inert gas as the sputter species (Ar, Kr, etc.), so as to not react with the

target material, most common of which is argon. For oxide growth a reactive sputter process is

commonly employed by introducing oxygen with argon to ensure the films are fully oxygenated.

The ratio of argon to oxygen can vary dramatically but typically includes some amount of argon.

Contrarily, oxide PLD growth often has no use for a process gas such as argon and as such only

varies the pressure of oxygen to change the plume dynamics and oxidation state.

While the effect of total partial pressure on the RHEED signal is trivially known, higher

pressures lead to an increased number of scattering events between the electron beam and the

gas, the effect of the gas species is low commonly considered. For this we consider the

difference between the differential scattering cross section between argon and oxygen. This

takes into account the Bohr radius for each element and the electron beam energy in order to

quantify the amount of scattering given a specific incidence angle. For this we can first consider

a fairly typical RHEED energy of 20 keV and then simulate23

and compare the scattering cross

section between argon and oxygen as shown in Figure 4.1. What is clearly evident is that the

differential scattering cross section for argon is approximately 4 times larger than that of oxygen

demonstrating that the amount of the electron beam that we will be scattered is 4 times larger and

this remains approximately true up to very large scattering angles. This has a major impact when

comparing sputtering with PLD with regards to the effect on the RHEED beam. From this

83

calculation we can expect to observe a much larger amount of diffuse scattering during

sputtering due to the presence of argon gas, increasing the overall background intensity.

From these simulations we can then begin to consider ways to reduce the overall

scattering caused by the gas. Since we are limited in the number of variables this is very

straightforward but none the less something that is important to consider and rarely discussed in

the literature related to high pressure RHEED. By increasing the electron energy from 20 keV to

35 keV which is the upper limit of our power supply, we can see a noticeable decrease in the

scattering cross section starting around 1 degree of theta or the incidence angle as shown in

Figure 4.1.23

The decrease in the number of scattering events that occur at greater than 1 degree

is significant and as such we will be work at 35 keV for the remained of the thesis. This increase

in the electron energy does have consequences which need to be taken in account as well. The

higher energy results in an enlarged Ewald’s sphere and thereby contracting the diffraction

pattern as observed on the phosphor screen. This may not be a negative impact depending on the

Figure 4.1. The differential scattering cross section for Ar and O atoms are shown as a

function of theta. The scattering cross section is shown for both 20 and 30 keV electron energies

for each atom type. It can be seen that Ar has a significantly larger scattering cross section than

oxygen and that higher electron energies help reduce the scattering at higher angles.23

84

overall RHEED setup but should be known as it can lead to a decrease in resolution, the number

of discernable pixels between diffraction spots.

In reality we can observe these calculated effects with the most noticeable change coming

when the gas species is varied. For this experiment the total pressure in the chamber was fixed at

75 mTorr, the RHEED power supply was also held constant at 35 keV including the filament

current, grid, and focus parameters. The only variable that was changed was the gas species in

the chamber so as to observe a direct comparison. This effect is seen in Figure 4.2 where the

change from an oxygen gas species to argon species is observed by capturing a diffraction

pattern from a SrTiO3 substrate near room temperature. Both images show a sharp diffraction

pattern typical of a high quality STO substrate. What is noticeably difference is the overall

brightness of the patterns. The pattern captured in the oxygen environment appears dimmer

when compared with the one from the argon environment. This can be understood as the

majority of the diffraction intensity is focused in the spots where the camera is saturated and very

little of the beam intensity is observed as diffuse scattering. Conversely, the image taken in the

Figure 4.2. Scattering comparison for Ar and O2 gas. A RHEED pattern of a STO substrate is

shown in a and b in oxygen and argon environments respectively. The diffuse background intensity

is noticeably larger in the argon environment as predicted by argons larger scattering cross section.

85

argon environment is very bright with measurable intensity in nearly every pixel shown. This is

a result of the higher scattering cross section of argon, scattering the electron beam in all

directions from the diffraction pattern.

The consequences of this effect are that the measurable dynamic range decreases due to

the effective minimum value increasing from 0 to a nonzero value. As an example, when

monitoring intensity oscillations, one can imagine a scenario where the actual minimum intensity

is a lower value than the background intensity in which case a significant amount of diffuse noise

would be measured in this situation rather than a smooth curve. Additionally, the background

intensity becomes more significant as the total pressure increases. Two possible remedies for

this issue are to increase the electron energy as previously mentioned but also to reduce the

working distance of the electron beam so as to limit the number of scattering events that occur.

For this work 35 keV was used and a total working distance from electron gun aperture to

phosphor screen of 20 cm although shorter distances are recommended keeping in mind that both

of these adjustments make a contraction in the pattern spacing.

4.2.2. Deflection due to magnetic field

The effect on the electron beam stemming from the scattering due to gas is significant

and deserves attention but the effect from stray magnetic fields is a much larger issue. This is

such an important issue that in fact an entire chapter of this thesis is dedicated to the handling of

the fields produced by the sputter sources. Knowing that a more in depth treatment of this topic

will be presented later on, in this section it will suffice to outline the premise of the issue and

ways to deal with it at a very basic level.

86

As with any instance where a focused electron beam is used, great care is taken to the

shield the apparatus from stray magnetic fields as they can deflect the beam in unpredictable

ways. In particular with regards to our RHEED setup, the electron gun is completely wrapped

with μ-metal to shield from undesired magnetic fields. μ-metal is a nickel-iron alloy that has

very high magnetic permeability and works by reshaping the magnetic field lines to that of shield

so that they go around the apparatus. In addition to the electron gun being magnetically shielded,

the extension tube inside the chamber is also shielded in μ-metal. In this way only the area

where the electron beam is exposed is it susceptible to external magnetic fields.

Typically the primary sources for stray magnetic fields are electromagnetic motors, DC

currents, permanent magnets (i.e. the full range vacuum gauge), and the geomagnetic field.

While these fields exist everywhere they are often manageable for RHEED systems and do not

pose a significant problem. The issue when discussing RHEED with sputtering is that

magnetron sputter sources by definition have very strong permanent magnets in them with

magnetic moments on the order of several thousand Gauss. The importance of these magnets,

one center cylindrical magnet and an outer annular magnet, for the sputter process was

previously discussed in Chapter 1. The crux of the issue is that in order to use RHEED during

sputtering these strong magnets must be placed in close proximity during 90° off-axis sputtering

to the section of the exposed electron beam which results in significant deflection of the beam

trajectory.

The degree to which the beam is deflected and the new trajectory path is dependent on

relative orientation and position of the sputter sources to the beam, the number of sources, the

sizes of the sputter sources, and the working distance of the electron beam. All of these degrees

87

of freedom make this a complex problem to solve in simple terms and as such finite element

modeling has been employed to model our system and the electron beam trajectory for several

common growth geometries. This is the focus of the following chapter where we present these

simulations and also discuss novel ways to mitigate the deflection.

For this chapter it is enough to acknowledge that this is a serious problem and that for the

in situ characterization discussed later in this chapter the magnetic field was dealt with through

the large amount of flexibility and degrees of freedom that were built into the growth chamber

originally for other purposes. A schematic of the growth chamber and deposition setup is seen in

Figure 4.3 which highlights the 90° off axis geometry, the relative RHEED arrangement and

some of the important degrees of freedom used to counter the deflection of the electron beam.

The general procedure for aligning the electron beam is as follows; we begin with the magnetron

sputter source pulled away from the electron beam as far as possible to avoid major deflections.

Then, as it is inserted into the chamber we are able to track the deflection of the electron beam

and correct for the deflection as best as possible while still maintaining a diffraction pattern.

This requires adjustments of azimuthal angle, tilt angle, and also repositioning of the electron

gun using the X-Y translation stage. As best as possible the XYZ heater manipulation stage is

not adjusted so as to maintain the same growth working distances. This procedure is done in

small steps and is a slow process taking upwards of an hour to complete. If the sputter source is

quickly introduced into the chamber to the final growth position the electron beam will bend so

severely that it will no longer be viewable on the screen at which point it is very challenging to

find again.

88

Using all of the degrees of freedom at our disposal we are able to keep the electron beam

on the very edge of the screen but we are at our limit for our standard growth position. Any

more significant bending and we would not be able to recover the pattern on the screen. Once in

the growth position we are still able to observe a diffraction pattern from an STO substrate but

the pattern is severely tilted and the diffraction spots are no longer sharp but are streaky. This is

demonstrated in Figure 4.4 where a comparison between a STO pattern without a field and one

with the magnetron sputter sources is shown. Additionally a RHEED image is shown at the

Figure 4.3. Schematic of the growth chamber demonstrating the relative orientation between the

heater and the sputter gun is shown along with the RHEED setup. Required degrees of freedom of

the heater and electron gun are shown as the tilt and azimuth and the x, y, and beam tilt, respectively.

89

growth pressure demonstrating the combination of the magnetic field and scattering due to gas.

However, this result is adequate because the most important factor is that we can observe and

monitor the pattern in particular the specular spot which we are able in this geometry. This

means that we are able to proceed with the growth of oxide materials while in situ monitoring

with RHEED.

4.3. RHEED intensity oscillations

Despite the fact that to this point we have not solved the issue of the beam deflection due

to the magnetron sputter sources we are able to observe a RHEED diffraction patter at the growth

pressure and with a single sputter source in the growth position as schematically shown in Figure

4.3. This means that the specular spot intensity can then be monitored in real time during the

sputter growth with the only a few remaining questions. First, how does the electric field

Figure 4.4. Sputter deposition effects on RHEED. A STO substrate in vacuum is shown in a with sharp

diffraction spots and Kikuchi lines. The effect from the magnetron sputter source on the diffraction

pattern is seen in b, where the pattern is clearly tilted the diffracted spots are less sharp and the Kikuchi

lines disappear. c shows the tilted RHEED pattern with the addition of 200 mTorr of Ar gas where a large

diffuse intensity obscures the pattern further.

90

produced by the sputter power supply and the plasma itself affect the electron beam and

observed intensity on the screen? Second, what is the growth mode of oxide films grown by 90°

off-axis sputtering? The later question will be addressed thoroughly in the remainder of this

chapter through the deposition of three separate well studied oxide materials; LAO, LSMO, and

SRO. The first two materials, LAO and LSMO, are known to grow in the layer by layer mode

during PLD growth with clear RHEED oscillations observed.24-26

SRO growth on the other

hand, has been well established to grow primarily in the step flow regime following the first few

unit cells.27,28

These materials will help address the way in which sputter deposition is similar or

different to PLD and if it possible to observe RHEED oscillations during sputter deposition.

However, before continuing on to the actual depositions and RHEED analysis, the first

question will be addressed. Typical of most oxide sputter growths, an RF power generator was

used for these growths to avoid charging on the target surface. Despite the oscillating potential

of the sputter source there is still a noticeable effect when the plasma is struck. It has not been

studied to a great degree but the intensity of the specular spot occasionally does not change, has a

sharp increase in intensity, or a sharp decrease. We believe this is a product of the electron beam

being shifted very slightly when the sputter voltage is turned on and the initial alignment of the

electron beam on the sample. Such that when the electron beam is shifted slightly as a result of

the electric field the incidence angle changes rapidly and depending on the initial alignment it

can increase or decrease the intensity. This does not however appear to be a significant effect on

the resulting intensity monitoring since it is a one-time event then stabilizes and is not a major

change in the overall intensity.

91

The sputter plasma itself is full of ionized elements and free electrons by definition.

Additionally, the plasma is relatively large from a 2” sputter source and the deposition of the

target material happens everywhere. This can have a serious negative impact on the phosphor

screen. If care is not taken to protect the screen from the sputter plasma it will not only light up

as a result of the free electrons striking the screen but it will also become coated rather quickly

and become unusable. To solve this problem an aluminum foil cone with a small aperture, ~1”

or less, was formed and placed over the screen. The aperture was placed as close as possible to

the substrate such that the electron beam could diffract off of the substrate and then travel freely

to the phosphor screen inside of the cone. This greatly reduced the line of sight deposition from

the sputter source and the overall interaction between plasma and the screen since the conduction

through the smaller aperture is poor. With the added cone, degradation of the phosphor screen

has not been noticed after tens of sputter depositions, while without the cone a change can be

noticed after the first deposition. This may not be the perfect solution but the point of this

discussion is to make it clear that the interaction between the phosphor screen and the plasma

cannot be ignored and this is the best solution we were able to find. With this problem solved we

can now turn to the monitoring of sputter deposition with real time RHEED.

4.3.1. LaAlO3/SrTiO3

92

As discussed in the previous chapter, the LAO/STO system is an ideal model

heterostructure for in situ RHEED with sputtering due to the high quality epitaxial growth that

we demonstrated by sputtering29

and the strong electrical dependence on the thickness.30

With

this in mind we chose this material system to be the first growth we monitored by RHEED with

the hope of observing layer-by-layer intensity oscillations. Due to the lack of real time analysis

techniques during sputtering, it was unclear what to expect in terms of growth mode.

In Figure 4.4c the starting RHEED pattern of the TiO2 terminated STO substrate is shown

at the growth conditions as previously discussed. The background intensity and tilting of the

pattern is clearly visible. Despite the poor signal to noise ratio the pattern is present and during

Figure 4.5. RHEED oscillations during a LaAlO3 growth on a SrTiO3 substrate showing

specular spot intensity oscillations that get weak very quickly but then are maintained out to greater

than 10 oscillations. The inset shows a zoomed in region to better display the oscillations with low

intensity. The large amount of noise is a result of the diffuse scattering from Ar gas screening the

low intensity RHEED pattern shown where the spots are barely observable above the background.

93

the deposition the specular spot and the diffracted spot intensities were monitored. The resulting

intensity versus time is shown in Figure 4.5, where intensity oscillations can be seen from both

the specular spot and diffracted spot out to greater than 15 oscillations. This clearly establishes

that RHEED oscillations can be observed during sputtering for the first time and that sputter

deposition can proceed in an extended layer-by-layer fashion, at the very least for LAO growth

on STO. The specular spot and diffracted spot intensities oscillate out of phase which is

commonly observed31

and they have a period of 49 seconds. The data is rather noisy and the

intensity dampens quickly after the first few oscillations. This dampening is fairly frequently

observed by other groups,30,32

suggesting that the sputtered LAO growth dynamics are

comparable to the PLD grown films. Following the dampening the intensity stabilizes and the

oscillation amplitude is just large enough to appear above the noise level. When considering the

static RHEED image shown in Figure 4.5 from the mid-zone where the intensity of the spots are

barely above the background noise it can be understood how the recorded amplitude would be

weak in comparison to the noise level. None the less, RHEED oscillations were reproducible

during sputter growth of LAO films confirming that RHEED with sputtering is promising as a

reliable in situ analysis technique that has been greatly needed during sputter deposition.

To confirm precisely that each of these RHEED oscillations are in fact stemming from

one unit cell of deposition in the layer-by-layer regime, an ex situ XRR measurement was

performed as seen in Figure 4.6a. From the fit performed of the XRR measurement an accurate

thickness of 11.4 nm can be obtained. Using this thickness, the total deposition time of 1500 s,

and the out of plane lattice parameter, the unit cell deposition rate can be calculated to compare

with the RHEED oscillation period. The out of plane 2θ – θ scan is shown in Figure 4.6b where

94

the (002) film and substrate peaks are shown. This gives an accurate out-of-plane lattice

parameter of 3.74 Å for the LAO film which is important for the calibration since this value can

vary as stoichiometry and the strain state change. The unit cell deposition rate obtained from the

thickness, time, and c - lattice parameter is 49 s/uc which matches closely with what was

obtained by the RHEED analysis confirming that the RHEED oscillations do in fact correspond

to one unit cell of growth. With this real time information it would then be possible to accurately

control the thickness of the film and thus the conductive state of the interface using RHEED.

Figure 4.6. Ex situ structural and surface analysis of the LAO/STO film. An XRR measurement

is shown in a with the corresponding simulation used to get an accurate film thickness. The out of

plane XRD measurement in b of the LAO film grown on STO indicates an c-lattice parameter of 3.74

Å which is used to form a growth rate calibration. The AFM image seen in c shows the LAO film has

smooth step and terrace structure similar to the substrate.

95

Ex situ analysis was further carried out on this film with AFM shown in Figure 4.6c. The

surface of the film is comparable to the surface shown in the previous chapter with an rms

roughness of 0.15 nm. Additionally, the rocking curve of the out of plane 002 peak is shown

with a FWHM of 0.08° which is also comparable to what was previously shown. This

information further supports what is determined with the in situ RHEED by the images taken

after the deposition which show clear sharp spots indicative of a high quality crystalline material

and a very smooth surface. This demonstrates that a lot of information about the film can be

obtained in situ without ever exposing the film to air and performing time consuming ex situ

measurements.

4.3.2. La0.7Sr0.3MnO3/SrTiO3

The RHEED data taken on the LAO films was a good starting point as it was the first

demonstration of the RHEED oscillations during a sputter grown film. However, the data was

very noisy throughout and the oscillations had small amplitudes and not well defined out to a

large number of periods. It was determined that a large portion of this noise was due to external

vibrations and noise entering the system by vibrating both the electron gun as well as the heater

mount among other places. To avoid this noise and improve the signal a number of steps were

taken to isolate the chamber from the external vibrations. Additionally, the sampling density of

the RHEED pattern was reduced and thereby increased the averaging time for each data point.

The capture rate is still fast enough to capture 3 frames per second so as to not miss any fine

96

details in the RHEED signal. By taking these measures the signal to noise ratio was significantly

improved.

To further prove the robustness of the RHEED with sputtering technique we chose to

grow LSMO films of STO substrates. This material system was chosen for several reasons; not

only is it a widely studied material system also with a critical thickness26

but also because it is

known to show strong RHEED oscillations when grown by PLD.25,26

In this way we could grow

a relevant material system that is known to grow layer by layer with PLD and also one that will

hopefully bring the signal level of the RHEED oscillations well above the nose level although

the scattering from the gas will still be present.

Figure 4.7. La0.7Sr0.3MnO3 growth on a SrTiO3 substrate showing clear RHEED

intensity oscillations indicating a layer-by-layer growth mode. The inset shows a

diffraction pattern of the film after growth in vacuum with the magnetic field present,

sharp spots with some streaking is seen indicating a flat 2-dimensional surface.

97

The LSMO films were grown at 750 °C in a 200 mTorr gas environment of 34:1 sccm of

Ar to O2 from a stoichiometric ceramic LSMO target. Films grown at 135 mTorr, with the same

gas ratio as above, also exhibited similar results (not shown here). The films were grown with an

RF power of 50 W in the 90° off-axis geometry. The fact that LSMO is a magnetic material, the

presence of the target further complicates the bending of the electron beam in addition to the

magnetron sputter source.

Despite this added field the initial RHEED pattern from the STO substrate matches

closely to what was previously shown with the clear tilting of the pattern. Monitoring the

specular spot intensity during the LSMO growth gives the resulting RHEED oscillations as can

be seen in Figure 4.7. The inset in Figure 4.7 shows a diffraction pattern of the LSMO film after

the growth in vacuum where sharp spots can be seen with some minor streaking indicating a very

flat 2-dimensional surface and good crystalline quality. The RHEED intensity data from the

(0,0) specular spot versus time for the LSMO growth is a clear improvement from the LAO

growth, with strong oscillations that decay with time which is very similar to previously reported

PLD growth of LSMO.26

Once again this suggests that the LSMO growth dynamics are very

similar when comparing PLD to sputtering. This also indicates that layer by layer growth by

sputtering is not specific to just one material system and is probably quite robust.

In order to characterize the structure and surface of the LSMO film grown on STO, x-

ray diffraction (XRD), x-ray reflectivity (XRR), and atomic force microscopy (AFM)

measurements were carried as seen in Figure 4.8. The out-of-plane XRD plot seen in Figure 4.8a

shows the substrate and film (002) peaks. The film peak has clear Kiessig Fringes indicating the

surface and interface are very smooth and sharp. The rocking curve of the LSMO film (002)

98

peak in Figure 4.8b has a FWHM value of 0.026° which is on the order of the STO substrate and

indicates high epitaxial crystalline quality.

In Figure 4.8c the XRR of the LSMO film on STO is seen with a fit which shows good

agreement. From the XRR data and the fit we can get a very accurate measurement of the film

thickness and also insights into the surface and interface quality. From the XRR thickness

calibration we find a growth rate of 32 s/uc. This matches very well with the RHEED

oscillations which also show an average period of 32 s/uc. However, a closer look shows that

Figure 4.8. X-ray and AFM measurements of the LSMO film. a) Out of plan 2θ-θ scan showing

the LSMO and substrate (002) peak with clear Kiessig fringes. b) Rocking curve of the LSMO (002)

film peak with a FWHM of 0.026°. c) XRR scan of the heterostructure used for thickness calibration

from the simulation. d) AFM image of the LSMO film showing clear step and terrace structure.

99

the period starts slightly dilated during the first 6 unit cells before finding its steady state growth

rate. The far less noisy data for the LSMO in comparison to the LAO data permitted us to obtain

this finer level of analysis. This demonstrates the power of RHEED to have careful layer by

layer control to a finer degree than can be achieved by a calibration and timed growth especially

when very thin films are desired. Figure 4.8d shows the AFM image of the film which clearly

shows the step and terrace structure preserved in the film from the treated STO substrate. This

confirms the smooth surface as expected from the RHEED pattern, the Kiessig Fringes and the

XRR fit.

4.3.3. SrRuO3/SrTiO3

The LSMO film growth confirmed that the sputter with RHEED technique is in fact quite

robust and likely applicable to any material system that can grow in the layer by layer regime.

To further test the robustness of this technique we turned to SRO growth which is commonly

believed to grow in a step flow regime.27,32

While intensity oscillations are a characteristic sign

of layer by layer growth, step flow growth has a RHEED signature of no oscillations and simply

a continuous bright intensity. The PLD grown SRO films typically start in the layer by layer

regime before transitioning to step flow after the first few unit cells.27

This makes the SRO

materials system a very interesting one to observe by RHEED during sputter deposition since

epitaxial SRO films were first deposited by sputtering7 and it is commonly assumed that they

also adopt the same growth mode as observed in PLD. This assumption is made due to the lack

of in situ analysis options available during sputtering. SRO is also an important material system

100

as it is commonly used as an epitaxial bottom electrode in oxide heterostructures due to SRO’s

excellent lattice match with many common substrates.

We grow SRO thin films on TiO2 terminated STO substrates in the 90° off-axis geometry

with in situ RHEED. The sputtered SRO films were grown at 590 °C with a working pressure of

200 mTorr. The gas ratio for this growth was 12:8 Ar to O2 respectively and the RF power was

fixed at 100 W. Using in situ RHEED we can identify both the growth mode and observe

RHEED oscillations during the deposition as seen in Figure 4.9a. The intensity of the (0,0)

specular spot was monitored as a function of time during the deposition. Figure 4.9a clearly

shows that the film grows in the layer by layer mode with oscillations that extend out to greater

than 50 clear intensity oscillations, each corresponding to a single unit cell of deposition (only

shown out to 40 here to be able to distinguish peaks). This allows for the ability to have exact

unit cell control of the growth during sputter deposition, which has never previously been

demonstrated for SRO growth out to more than ~10 unit cells and never during sputter growth.

Figure 4.9. RHEED oscillations from SRO growth are seen in a) as the black line. The oscillations

are shown out to 40 unit cells with the red line showing the fit. The inset highlights the first few unit

cells where the extended period can be seen. b) is a plot of the period versus unit cell displaying the

relaxation time required to reach the approximate steady state with an exponential decay fit in red.

101

Furthermore, the extended oscillations are of particular interest when comparing the

sputter grown SRO to that which is observed in PLD grown SRO. In PLD grown films, it has

been shown that the growth mode of SRO begins as layer by layer growth and transitions to step

flow growth after several unit cells.27

In addition, these PLD grown films show a RHEED

signature corresponding to the transition from RuO2 termination to SrO termination.33

We find

that this change in termination can also be observed in sputter grown SRO films, as seen in

Figure 4.9a and made clearer in the inset. The elongated first unit cell oscillation of

approximately 38 seconds is nearly double the steady state average of 19.3 seconds per unit cell

and corresponds to the transition from BO2 termination to AO termination. However, in contrast

to PLD growth, the SRO does not transition from layer by layer to step flow growth. Figure 4.9a

shows RHEED oscillations characteristic of layer by layer growth with each oscillation

corresponding to one unit cell. If the growth mode were to have transitioned to step flow, the

oscillations would have died out. The oscillations were observed out to greater than 50 unit

cells, clearly establishing the ability to have layer by layer control over sputter grown SRO films.

Due to the transition in growth mode of PLD grown SRO, layer-by-layer control out to many

unit cells has not been demonstrated, giving sputtering better thickness control over SRO films.

To further demonstrate the capability of this technique, these films were grown without

presputtering in order to observe the time necessary to reach steady state. A fit of the oscillations

is shown in red in Figure 4.9a starting with the first complete peak. This fit was used to obtain

the crest to crest period of each oscillation, or the time required for each successive unit cell,

which is plotted in Figure 4.9b. From this plot it can be observed that unit cell growth rate starts

near 21 seconds per unit cell and relaxes down to the steady state period of approximately 19.3

102

seconds after 15 unit cells following approximately an exponential decay shown by the red

curve. From this we can determine the minimum presputter time but can also observe that it is a

non-uniform growth rate throughout, indicating that without the use of an in situ monitoring

technique precise unit cell control may not be possible.

Figure 4.10. Topographic and structural characterization of the SRO film by AFM, XRD, and

XRR. (a) AFM image of the SRO surface showing single unit cell steps and incomplete step edges.

The inset shows the RHEED image of the SRO with sharp spots after the growth in vacuum with the

sputter source in position causing the tilting of the pattern. (b) Out of plane XRD scan showing the

relationship between the (002) peak of the SRO and STO. Distinct Kiessig fringes are present

indicating the interface and surface are both sharp and smooth. (c) The XRR data shown in black with

the fitted data in red provide an accurate thickness estimation, and indicate the high quality of the

surface and interface. The fit gives an estimated total thickness of 27.1nm.

103

The RHEED image of the 27.1 nm SRO film after the growth is seen in the inset of

Figure 4.10a. This image is taken in vacuum with the magnetic field from the sputter source still

present, which is evident by the tilted pattern. Clear diffracted and specular spots can be

observed with minimal streaking suggesting high crystallinity and a predominantly two-

dimensional surface. Ex situ analysis of the SRO film structure and thickness is carried out with

x-ray diffraction (XRD) and x-ray reflectivity (XRR) respectively and the surface morphology is

characterized by atomic force microscopy (AFM), as seen in Figure 4.10. The AFM image in

Figure 4.10a shows clear terrace structures; however, the steps do not have linear and smooth

edges, and show some small single unit cell islands on the terraces. This is comparable to SRO

films grown on STO by PLD,33

but the existence of the small islands on the sputter grown film

may be an indication of nucleation sites corresponding to the layer by layer growth as opposed to

the step flow growth of PLD. The AFM result also corroborates observations from the RHEED

pattern that the film surface is very smooth.

The out-of-plane XRD scan seen in Figure 4.10b shows the relationship of the (002) SRO

film peak to the STO (002) substrate peak. Clear Kiessig Fringes can be seen, which indicate a

smooth surface and interface. The rocking curve of the (002) SRO peak has a FWHM of 0.024°

indicating low mosaic spread and high crystalline quality. These ex situ measurements confirm

what has already been observed in situ with RHEED, demonstrating the power and efficiency of

real time monitoring.

The XRR data and corresponding fit can be seen in Figure 4.10c with the raw data in

black and the fitted data in red. Using the fit, an accurate total thickness of the film is found to

be 27.1 nm from which a growth rate of 20.1 seconds per unit cell is derived. When compared to

104

the 19.3 seconds per unit cell steady state rate acquired from the RHEED oscillations it can be

seen that these results are in close agreement. However, from the varying period of the

oscillations seen in Figure 4.9b it is clear that using RHEED is a more robust and reliable way of

ensuring precise unit cell control compared with a time based calibration. Additionally, these

results confirm that the SRO grows in a layer by layer mode because one would expect these

values to be significantly different if there was also a step flow contribution to the growth. This

is a clear indication that each RHEED oscillation corresponds to exactly one unit cell being

deposited on the sample surface, and as such, is the first demonstration by any growth method of

greater than 50 RHEED oscillations providing digital control during SRO growth.

4.4. Conclusions

This chapter clearly establishes the difficulty associated with RHEED during sputtering

and likely the reason to why the use of RHEED as a real time monitor is not more prevalently

observed during sputter deposition. The scattering due to the gas has been clearly demonstrated

for both oxygen and argon gas and ways to diminish the scattering effect have been laid out. For

this chapter it was sufficient to acknowledge the magnetic field existed and ways to deal with it

through flexibility without actually solving the issue. The magnetic field issue will be further

addressed in the following chapter through simulations of the field and the effect it has on the

electron beam. Despite the challenges we demonstrate that the RHEED pattern from the

substrate can be observed in the 90° off-axis sputter conditions. It naturally follows then that the

intensity versus time can be recorded from the pattern during sputter deposition which was

105

demonstrated in this chapter for LAO, LSMO, and SRO growth. All three of these scientifically

relevant materials demonstrate layer by layer growth which is of particular interest for the SRO

growth. With these results we can clearly say that digital control of sputter growth by in situ

RHEED is possible creating a new avenue for real time analysis during sputtering which

otherwise did not readily exist.

4.5. Reference for chapter 4

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5. Finite element modeling

5.1. Introduction

Throughout this thesis the theme of combining RHEED with sputter deposition is ever

present. Along with this theme is the notion that one of the limiting factors for combining these

processes together is the presence of strong permanent magnets in the sputter sources.

Previously, groups have acknowledged that this is an issue but since they had such a large

working distance it was not a significant factor.1 In our own work we have also experienced this

issue but were able to accommodate the bending of the electron beam through large degrees of

freedom which are not always sufficient. One easy solution may be to sputter without the

magnets which is possible yet extremely uncommon. However, the purpose of this chapter is to

not only understand the effect of the magnets on the electron beam but to find a robust solution

where RHEED with magnetron sputtering can be utilized in a large range of growth geometries.

To this point we have demonstrated that not only is RHEED an extremely powerful and

useful in situ analysis tool but also that critical information can be obtained from monitoring the

RHEED pattern during sputtering, such as unit cell oscillations. In order to take this one step

further and predict the effect of the magnetic field on the electron beam trajectory, COMSOL

(COMSOL Multiphysics® version 5, AC/DC and Particle Tracing Modules) was used to create a

finite element model of the magnetic field produced by the magnets in a 2-inch planar magnetron

111

sputter gun oriented in the 90° off-axis geometry, and to simulate the beam trajectory when

passing through the magnetic field. The finite element method is a useful strategy that solves

complex partial differential equations which collectively describe the entire system.2 The domain

of the problem is divided into a group of subdomains with each subdomain represented by a set

of element equations that locally approximate the original equations for the problem.

Recombining these local solutions into a single global solution yields accurate results for the

solution of the problem.

Finite element analysis (FEA) is a computational tool which utilizes the finite element

method. FEA includes the use of meshes which are generated in order to divide a complex

problem into small elements. Generally, more complex areas require finer meshes resulting in

smaller local elements and a larger amount of computing power. Once the object of interest is

discretized, the correct material properties and boundary conditions are applied to each small

element. These small elements are then used to solve the global problem. COMSOL

Multiphysics is a powerful tool capable of simulating complex physical systems. COMSOL uses

FEA and partial differential equations techniques to accurately simulate physical systems input

by the user.

The core COMSOL program enables simulations of mechanics, and additional modules

are available to expand the range of possible simulations. For this work, COMSOL core as well

as the ACDC modules and particle tracing module were used. The ACDC module allowed for

the simulation of a magnetic field produced by permanent magnets and the particle tracing

module allowed for the simulation of an electron beam as it moves through the magnetic field.

112

5.2. Modeling parameters

The object geometries were created in SolidWorks to scale and then imported into

COMSOL. The electron gun, phosphor screen, and heater were simulated as Type 316 steel, with

relative permeability of 1. The magnets were modeled with a relative permeability of 1.05 based

on NdFeB rare earth magnets and a remnant flux density of 1.201 T. All of the remaining

volume was simulated as a low pressure gas with relative permeability of 1. The particles in the

beam were given the mass and charge of an electron, and 25 particles were released from a 1 mm

wide aperture (exaggerated from actual dimensions slightly to make more obvious) with a

velocity of 1.05 x 108 m/s corresponding to 35 keV of kinetic energy.

The particle beam experienced a magnetic force from the magnets once the particles were

released. A mesh was created for the entire geometry and then a time dependent study was

performed. Post processing included coloring the north and south ends of each magnet red and

blue respectively as well as all the objects to easily distinguish them. Three-dimensional images

were produced showing the beam path via colors corresponding to the magnetic force exerted on

the beam at each point. A two-dimensional vector field showing the magnetic field strength and

direction was also created. After post processing, camera views were chosen to export the proper

images for use as figures.

5.3. Single and two gun 90° off-axis geometries

113

We simulated many magnetic orientations within several different geometries in order to

predict the bending of the beam and to find an optimal geometry for minimal beam deflection

caused by the strong magnetic fields present. The ideal magnetic field for our purposes would be

parallel with the electron beam’s direction which would result in zero deflection. The next best

solution is one in which the field strength along the beam path is reduced and the field direction

is constant for the entirety of the trajectory. This would reduce the magnitude of the deflection

and also make the deflection more predictable as it would be uniaxial. In particular, magnetic

fields that are perpendicular to the electron beam and horizontal will push the beam only

vertically, still allowing much control to overcome the problems caused by the deflection of the

beam. This concept is the basis of our findings and will be elucidated more clearly in this

section.

The first model created was representative of the growth geometry used for the previous

growths, one sputter gun in the 90° off-axis geometry, such that a better understanding of the

deflection we observed could be obtained. Subsequently we modelled a two gun 90° off-axis

geometry with facing guns to enable the growth of heterostructures and ideally balance the field.

A 2-dimensional cross-sectional view of the simulated magnetic vector field distributions can be

seen in Figure 5.1. The cross-section is a slice through a plane containing the heater and sputter

sources, with the electron gun positioned out of the plane of the page and the phosphor screen

into the plane of the page such that the electrons travel into the page. All of the working

distances used for these simulations are to-scale representations of our actual chamber setup.

Figure 5.1a shows the magnetic field close to the sample for the simplest case of a single gun in

the off-axis geometry. Here we can note the relative field strength and additionally the direction

114

of the field lines close to the sample. In the single gun setup the field lines point close to vertical

with some tilt toward the magnets.

If we consider an electron beam traveling straight into the page at the point directly above

the center of the sample in Figure 5.1a, the Lorentz force will point perpendicular to the

magnetic field direction toward the magnets with a magnitude proportional to the magnetic field

strength. Consequently, the beam will bend laterally in the x-direction towards the magnets at

this point. Similarly, COMSOL can be used to explore more complex chamber geometries. In

Figure 5.1b, a two-gun off-axis setup is shown with two sets of symmetrically facing magnets

(“symmetric” setup). In the cross section, an increase in the field strength near the sample is

observed along with a significantly transformed overall vector field distribution. Despite these

Figure 5.1. 2-dimensional cross-sections showing the magnetic field close to the sample for the single

gun, two gun symmetric and two gun antisymmetric magnet polarities. The electron beam direction

is into the page, and the RHEED phosphor screen is behind this plane (into the page). (a) Single gun off-

axis geometry shows the magnetic field lines close to the sample are pointing nearly vertically resulting in

a predominantly lateral Lorentz force. (b) Two gun symmetric off-axis geometry with the field lines

close to the sample completely vertical. (c) Two gun antisymmetric geometry displays the magnetic field

lines close to the sample are horizontal resulting in a Lorentz force in the y-direction. The magnitude of

the magnetic field can also be compared between the three setups.

115

considerable changes, the Lorentz force near the sample has similar results to that of the single

gun simulation with the force pointing completely laterally in the x-direction, now with a larger

magnitude.

In Figure 5.1c, the same two-gun facing geometry is shown as in Figure 5.1b except in

this case the magnet orientation of one gun is inverted such that the two guns are now in an

“antisymmetric” setup. In this setup, once again a change in the magnetic field lines is seen as

well but in this geometry the strength of the field near the sample is reduced. Most notably, the

magnetic field direction at the sample now points horizontally and the resulting Lorentz force

near the sample points downward in the y-direction. This significant change in the magnitude

and direction from both the single gun geometry and the “symmetric” setup renders the

“antisymmetric” geometry more desirable from a practical point of view (see discussion below).

The 2-dimensional cross sections of the magnetic field distribution shown in Figure 5.1

are a good starting point for understanding the field strength and orientation near the sample

surface and qualitatively predicting the electron beam deflection in different chamber

geometries. However, the 2-dimesional plots make the assumption that the beam is traveling

straight into the plane and do not consider the field in 3-dimensions. To quantitatively predict

the electron beam path a time-resolved finite element simulation is required, as provided by the

combination of the AC/DC and Particle Tracing modules in COMSOL. Simulations can be made

in COMSOL that use actual working distances between the electron gun, phosphor screen, heater

block, and magnets in the sputter source in our chamber geometry, including the 35 kV beam

voltage used experimentally, and can also simulate hypothetical setups. This provides a useful

platform for understanding what is observed and also for predicting and designing future

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systems.

Figure 5.2. Finite element simulations showing electron beam deflection in single gun, symmetric,

and antisymmetric sputter source geometries and their resulting RHEED patterns. (a) Single off-

axis sputter source layout with from a top view showing a clear deflection in the x direction is observed

with the beam missing the edge of the phosphor screen. (b) Two opposite facing sputter guns in the

symmetric geometry from a top view showing a more dramatic bending of the beam than in (a). (c)

Antisymmetric geometry is shown resulting in moderate bending only in the y direction and striking the

phosphor screen. The RHEED pattern seen in (d) is the image of a bare STO substrate in vacuum with a

single sputter source in the off-axis position. Clear tilting of the pattern is observed as well as a reduction

in the sharpness of the diffracted spots. (e) A RHEED image from an STO substrate is shown for the

antisymmetric setup exhibiting no tilting of the pattern and sharp specular and diffracted spots.

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Figure 5.2a, b, and c show 3-dimensional representations and simulations of the single

gun, symmetric, and antisymmetric setups, respectively. Side views of these simulations can be

seen in Figure 5.3. The color gradient along the beam shows the magnitude of the magnetic field

that the electrons encounter traveling from the gun to the screen. From these simulations the

electron beam path can be predicted for all three cases, providing a more complete understanding

of the influence of the magnetic field along the entire trajectory than the cross sectional views.

Figure 5.2a shows that the deflection due to the magnetic field from a single gun is sufficient for

the beam to completely miss the 55 mm diameter phosphor screen. It can be noted the Lorentz

force from the 2-dimensional cross-section (Figure 5.1a) predicts the beam to bend not only

towards the magnets but also slightly towards the heater. However, following the full trajectory

in 3D the net result is that the beam is actually bent away from the heater.

Figure 5.2b shows the 3D beam trajectory in the symmetric gun setup resulting in a

significantly larger deflection of the beam than observed for the single gun setup. From this, it

can be concluded that using the symmetric gun setup in our chamber would make it more

difficult both to have the beam diffract off the sample and have the resulting pattern strike the

phosphor screen. The large lateral deflections seen in the single gun or symmetric gun

geometries can be eliminated by inverting the magnetic polarity of one gun, as seen in Figure

5.2c for the antisymmetric setup. This geometry has the additional advantage of minimizing the

strength of the magnetic field along the electron beam trajectory (note the predominantly blue

color of the beam and the field strengths indicated by the color scale), resulting in dramatically

lower overall deflection of the beam. Crucially, because the field lines are parallel to the x-

direction along the entire beam path, the beam is only deflected in the y-direction and can easily

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be made to strike the phosphor screen. Upon a careful analysis of the 3D simulations, the field

lines change direction twice along the beam path leading to two zero-field points where the beam

bends first up and then down again.

Figure 5.3. Finite element simulations side view showing electron beam deflection in (a)

single sputter source, (b) symmetric dual, and (c) antisymmetric dual sputter source

geometries. (a) The single gun geometry is shown where the beam deflects above the surface

of the sample and also misses the screen. In (b), the symmetric setup is simulated which

shows a significant deflection in the -y direction, also resulting in the beam missing the sample

surface. The symmetric setup can be seen in (c) which shows a small bump in the beam

trajectory above the sample but the beam returns to the y=0 starting value at the screen.

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The presence of two sputter guns in the antisymmetric configuration has additional

benefits beyond permitting the use of RHEED during growth. First, the growth of complex

heterostructures or superlattices of two different materials would require in any case the presence

of two sputter sources. Second, it has been shown that an antisymmetric sputter geometry

improves film uniformity over large areas by eliminating the confined magnetic field created by

a symmetric setup that leads to resputtering of deposited materials; this makes scalability more

feasible.3 These factors all point to the 2-gun antisymmetric geometry as being ideal for

sputtering of complex oxides with in situ RHEED.

Experimentally the results match well with what is predicted from the simulations.

Without making any changes to the system, the direct beam in both the single gun and symmetric

setups is not observed where as in the antisymmetric setup the direct beam strikes the screen.

However, using all of the degrees of freedom in our chamber the beam can be redirected for the

single gun setup such that the beam diffracts off the SrTiO3 (STO) substrate and its diffraction

pattern can be observed on the screen. This pattern is shown in Figure 5.2d where the field

causes the diffraction pattern to noticeably tilt, and the spots become irregularly shaped and the

diffracted spots become very weak. Furthermore, the Kikuchi lines typically seen on a bare

substrate are no longer observable. For the symmetric case, the deflection is so severe that the

degrees of freedom are not enough to redirect the beam onto the screen and no diffraction pattern

is observable. For the antisymmetric setup, very minor adjustments were required in order to

observe a pattern from an STO substrate, shown in Figure 5.2e. There is no tilting present in this

pattern and sharp diffraction and specular spots are observed. The pattern does show some

minor streaking as a result of the magnetic field but strong Kikuchi lines are still present and the

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pattern is significantly more comparable to a typical STO pattern in the absence of a magnetic

field. This confirms the simulation results that the antisymmetric configuration is the ideal setup

for in situ sputtering with RHEED.

5.4. Two antisymmetric configurations

While the antisymmetric geometry is clearly preferable, further analysis is required to

fully understand the effects of the field on the RHEED pattern. Specifically, we should consider

the fact that there can be two possible configurations of the magnets; configuration A where the

field near the sample points to the left and configuration B where the field points the opposite

direction. This can be achieved by flipping the polarity of all the magnets such that we still have

antisymmetric magnets and the field shape is the same but the direction of the field points in is

opposite. This is shown more clearly in Figure 5.4 where both configurations are displayed. The

obvious change in impact on the electron beam between A and B is the direction of the Lorentz

force is opposite for each. However, what is not immediately obvious is what happens when the

beam is not traveling directly between the sputter sources.

To this point we have been considering an idealized setup where the sputter sources are

perfectly symmetric about the sample and the RHEED setup, and the electron beam travels

perfectly straight between the sputter guns. In reality this can be extremely challenging to

achieve and as such it is worthwhile to consider what happens when the system is varied. At a

basic two-dimensional level, as seen in Figure 5.4a and b, one can consider the Lorentz force not

only directly above the sample but also shifted horizontally for configurations A and B. It can be

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seen that for configuration A the Lorentz force always points back toward the center regardless

of which side of center the beam starts. Conversely for configuration B the Lorentz force always

has a component that points away from center.

In general this indicates that for configuration A the beam will always be re-centered and

for B the beam will always travel away from center unless starting perfectly centered. To get the

entire picture however, we must once again consider the full three-dimensional view which is

shown in Figure 5.4b. The three-dimensional view does in fact confirm what was originally

assumed as clearly shown by the simulations where the beam starts off-set in both positive and

negative x for both configurations. For A, the beam consistently is forced back to the center of

Figure 5.4. Magnetic field simulations for the two possible antisymmetric configurations are show in

a and b. Configuration A is shown in a while configuration B is shown in b. This effect is further

illustrated in the 3-dimensional models of the beam trajectory where an off-set in the x direction is

deflected back towards the middle for configuration A as seen in c and e, and for configuration B the

beam deflects away from the center in d and f.

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the screen whereas in B the beam is push away by the Lorentz force. This can have significant

implications when attempting to align the beam and additionally when any off-sets may exist.

Furthermore, at a finer level when only considering the RHEED pattern, configuration A actually

acts to elongate the pattern in the y direction whereas B will force the pattern to seem more

diffuse. Additional studies are needed to further understand the implications of these changes

and to create a more concrete conclusion about which configuration is more desirable, but for

now understanding that the difference exists is important.

5.5. Generic solution

The antisymmetric solution that was found in the previous section for 90° off-axis

sputtering clearly limits the magnitude of the deflection and also restricts the deflection to a

single axis making it more predictable and thus easier to account for. However, the 90° off-axis

geometry is a fairly specific growth geometry that is not necessarily applicable to all materials

systems. Being able to form a more generic rule of thumb that encompasses a much broader

range of growth geometries would therefore be significantly more beneficial for the wide-

ranging sputter community. Additionally, finding a setup where there is virtually zero effect

from the sputter sources would be extremely beneficial.

Using what was previously found in the two gun 90° off-axis setup, additional growth

geometries were simulated. What was found is presented in Figure 5.5, which shows a standard

two gun 45° confocal geometry as well as a setup that includes two sputter guns facing the heater

with the sample between them, occasionally referred to as a high rate off-axis setup.

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What can clearly be seen in both of these setups in that near sample surface the field lines are

horizontal when the magnets are antisymmetric. While the field strengths here are higher than in

the 90° off-axis setup, the same uniaxial bending as seen previously can be expected. Regardless

of the magnitude of deflection, the uniaxial bending is highly predictable and offsets can easily

be built into a system. It should also be noted that the working distances here are somewhat

random and the magnitude of the magnetic field will understandably scale with the proximity of

the sputter sources to the sample. From these simulations a generic guideline can be formed

which can be applied to a large range of growth geometries. Essentially, as long as two sputter

sources are present in some non on-axis yet geometrically symmetric setup with antisymmetric

magnet arrangement, a uniaxial bending perpendicular to the sample surface can be expected.

This expands the robustness of this technique to include nearly every angle of off-axis growth

and also every angle of confocal growth.

Figure 5.5. Confocal and high rate off-axis antisymmetric setups. Two sputter arrangements are

shown with antisymmetric magnet arrangements for 45° confocal sputtering in a and two gun high rate

off-axis sputtering in b. In both cases the field lines are horizontal above the sample resulting in a

uniaxial deflection of the electron beam.

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When considering pure on-axis growth, it is desirable to have a solution for this

arrangement as well. Even in on-axis growth, there is a need to balance the field in some way.

At first we considered an antisymmetric sputter source behind the heater facing the main source.

However, this resulted in a non-uniaxial and severe bending of the electron beam. Contrary to

what was previously observed, a symmetric setup was simulated and the two-dimensional

magnetic field cross-section is shown in Figure 5.6. For this model, it is more helpful to consider

the electron beam traveling horizontally directly above the sample. In this case all of the field

lines point horizontally as well and switch orientations at the center point. This results in a net

deflection due to the magnetic field of zero since the resulting Lorentz force is pointing in the

same direction or opposite direction of the beam path.

Figure 5.6. An on-axis sputter arrangement

with symmetric sputter sources centered

around the substrate surface is shown. The

field lines are all perfectly horizontal in this

plane indicating that an electron beam

traveling horizontally would no experience a

Lorentz force.

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Although it may not be an ideal setup to have another sputter source (or simply another

set of magnets) behind the heater and in certain situations not possible, it is the best arrangement

for completely avoiding deflection due to the sputter source. Furthermore, it is the only setup in

which field lines are parallel to the electron beam that we have found thus far. In addition to

this, one can imagine a situation where the heater can then rotate from one sputter source to the

other to deposit multi-target heterostructures while still incorporating RHEED. This is

reminiscent of the motions described for the hybrid chamber in Chapter 2 except in theory there

would be zero effect from the sputter sources. This setup has not yet been experimentally

demonstrated but it should be noted that a small variation of the electron beam from the central

axis will result in some bending of the beam. This is true for all of our simulations and it is also

worth mentioning that RHEED does require some finite incident angle of the beam on the

sample whereas here we are simulating the beam traveling parallel to the substrate surface.

Nonetheless this clearly demonstrates that not only are we able to form a general rule of thumb

for non on-axis setups but even for on-axis deposition the deflection can be completely avoided.

5.6. Antisymmetric SRO growth

To further demonstrate the power of the antisymmetric setup, SRO thin films were grown

using configuration A. While our current growth chamber does not have the ports needed to

allow the inclusion of two sputter source in the antisymmetric arrangement, we were able to add

a dummy set of magnets to act as if a sputter source was present. For this, a stand was

constructed to hold a set of sputter magnets perfectly opposite the functioning sputter gun with

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antisymmetric polarities. Oscillations of the RHEED specular spot intensity during the

deposition are readily observable, as shown in Figure 5.7a, and permit the calibration of the

deposition of single unit cells and allow the identification of layer-by-layer growth mode, as

discussed below and previously demonstrated. A Gaussian fit of the oscillations (the red line in

Figure 5.7a) was used to obtain the crest-to-crest period of each oscillation, and a plot of these

periods is shown in Figure 5.7b. After the initial few unit cells, the growth rate is fairly constant

throughout with an average period of 20.7 seconds.

Similar to the previously discussed SRO growth, the STO substrates are treated to form

a TiO2 termination4 prior to growth, and the corresponding atomic force microscopy (AFM)

image exhibits a single unit-cell step-and-terrace structure with nearly straight step edges is

shown in Figure 5.7c. The inset to Figure 5.7c shows the RHEED image of the STO substrate

prior to growth in vacuum; the sharp diffracted and specular spots confirm the high quality

crystalline substrate and smooth surface. The inset in Figure 5.7d shows the RHEED images

taken in vacuum after the SRO growth. The diffraction pattern still shows sharp spots with some

minimal streaking suggesting a high degree of crystallinity in the film and a predominantly two-

dimensional surface. The AFM image in Figure 5.7d shows a clear step-and-terrace structure

with single unit-cell step height and some small single unit cell islands on the terraces. This

morphology is comparable to SRO films grown on STO by PLD,5 but the existence of the small

islands on the sputter grown film may be an indication of nucleation sites corresponding to the

layer by layer growth deduced from the extended RHEED specular spot oscillations. The AFM

result also confirms observations from the RHEED pattern that the film surface is atomically

smooth.

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The film thickness determined by x-ray reflectivity (XRR) scan confirms that each

RHEED oscillation corresponds to the deposition of a single unit cell as shown in Fig. 4e. The

fit suggests that the surface and interface are smooth and the density is uniform throughout, and

yields a total film thickness of 10.9 nm corresponding to a growth rate of 21.7 seconds per unit

cell. This rate agrees well with the 20.7 seconds per unit cell steady state rate acquired from the

RHEED oscillations in Fig.4b. This confirmation has several implications. First, observation of

RHEED oscillations allows for an in situ, real time capability to control sputter deposition at the

level of a single unit cell. Second, extended RHEED oscillations identify a layer-by-layer

growth mode for SRO by sputtering, unlike step-flow growth mode that characterizes SRO

deposition by PLD after the first few unit cells.6 Third, analysis of the RHEED oscillations at

Figure 5.7. RHEED intensity oscillations during SRO growth and unit cell by unit cell growth rate

plot along with the corresponding AFM and RHEED images. The XRR data for the SRO film is also

shown. (a) Specular spot RHEED oscillations during SRO growth on STO. The actual data is seen in

black with the Gaussian fit in red. Clear oscillations are observed corresponding to one unit cell of

growth. (b) The peak to peak period for each unit cell of growth is shown. The inset highlights the first

few oscillations where the extended period can be seen due to the termination conversion during SRO

growth. (c) AFM image of the STO substrate prior to growth showing smooth step and terrace structure.

The inset shows the RHEED image of the STO substrate in vacuum before growth. (d) AFM image of

the SRO surface showing single unit cell steps and incomplete step edges. The inset shows the RHEED

image of the SRO film with sharp spots after the growth in vacuum and no tilting of the pattern. (e) The

XRR data (black line) and fits (red line) provide an accurate thickness estimation of 10.9 nm, and indicate

the high quality of the surface and interface.

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the start of growth allows the identification of a surface termination inversion in the first few unit

cells of deposition. As shown more clearly in the inset of Figure 5.7b as the shoulder to the

second peak, the period from the first peak to the shoulder is 1.5 times the steady-state period.

The termination inversion from RuO2 to SrO in the first few unit cells reported here for

sputtering is similar to that reported for SRO growth by PLD.5 However, in contrast to PLD

growth, the SRO does not transition from layer by layer to step flow growth after the termination

inversion. The extended RHEED oscillations that characterize layer-by-layer growth by

sputtering allows real time calibration and control of film thickness to the unit-cell level, as

compared to the time-based calibration required for the step-flow growth mode by PLD.

While these results are very similar to what was previously demonstrated, it should also

be noted that the oscillation amplitude is not as large during the antisymmetric growth as

previously shown for the single gun. Additionally, the oscillations do not extend out to as many

periods as in the single gun setup. This may be a result of the antisymmetric arrangement

affecting the dynamics of the plasma near to the sample surface similar to what Newman et al.

observed.3 This is further supported by the fact that the growth rate is slower for the

antisymmetric setup despite all other parameters being kept constant. Moreover, since there is

no evidence of the film surface roughening by AFM, it could be plausible that the growth is

converting to step flow growth in this regime more readily than observed in the single gun setup.

Nonetheless, this growth confirms that the antisymmetric setup produces an improved RHEED

pattern and also can be used to have layer by layer control during sputtered film growth.

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5.7. Conclusions

Finite element models were created for a wide range of growth geometries such that the

deflection of the electron beam due to the magnetron sputter sources could be reduced and

mitigated. For the single gun 90º off-axis sputtering setup that was used to grow films in the

previous chapter, the FEA was able to predict a similar bending to what was observed via

experiment. While this is useful for understanding the deflection and even designing setups that

can accommodate this bending the ultimate goal was to determine a way to mitigate or avoid all

deflections. This was accomplished through the addition of a second sputter source which

should have an antisymmetric magnet arrangement with respect to the primary sputter gun for

any non on-axis growth setups. This results in a uniaxial bending that his highly predictable and

greatly reduced from the single gun setup. For on-axis deposition, a magnetically symmetric

sputter gun facing the primary gun behind the heater will result in zero deflection of the electron

beam.

Further work is required to fully understand the effect of the two possible magnetic

configurations in the antisymmetric setup, but clearly there is a difference. Of significant interest

is the difference between the SRO grown via one gun versus the two gun setup. Nominally these

two growths should be the same but clearly the magnetic fields have an effect on the deposition

and this will also require further investigation to get a better grasp on what is happening. Now

that a basic guideline has been established for the inclusion of magnetron sputter deposition with

RHEED this should open the door for many interesting future studies.

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5.8. References for chapter 5

1 Svedberg, E. B., Birch, J., Edvardsson, C. N. L. & Sundgren, J. E. Real time

measurements of surface growth evolution in magnetron sputtered single crystal Mo/V

superlattices using in situ reflection high energy electron diffraction analysis. Surface

Science 431, 16-25, doi:10.1016/S0039-6028(99)00498-7 (1999).

2 Strang, G. & Fix, G. J. An analysis of the finite element method. Vol. 212 (Prentice-Hall

Englewood Cliffs, NJ, 1973).

3 Newman, N., Cole, B. F., Garrison, S. M., Char, K. & Taber, R. C. Double gun off-axis

sputtering of large area YBa2Cu3O7- superconducting films for microwave applications.

IEEE Transactions on Magnetics 27, 1276-1279, doi:10.1109/20.133417 (1991).

4 Kawasaki, M. et al. Atomic Control of the SrTiO3 Crystal Surface. Science 266, 1540-

1542, doi:10.1126/science.266.5190.1540 (1994).

5 Rijnders, G., Blank, D. H. A., Choi, J. & Eom, C. B. Enhanced surface diffusion through

termination conversion during epitaxial SrRuO3 growth. Applied Physics Letters 84, 505-

507, doi:10.1063/1.1640472 (2004).

6 Choi, J., Eom, C. B., Rijnders, G., Rogalla, H. & Blank, D. H. A. Growth mode transition

from layer by layer to step flow during the growth of heteroepitaxial SrRuO3 on (001)

SrTiO3. Applied Physics Letters 79, 1447-1449, doi:10.1063/1.1389837 (2001).

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6. Summary and future thoughts

In this thesis, a clear goal was set upon from the start with Richard Feynman’s prediction

guiding the way. The goal of using reflection high energy electron diffraction as a means to

garner atomic level control over film growth was the focus. Although RHEED is a well-

established technique, we attempted to push the commonly accepted bounds and limits to expand

RHEED into the realm of sputter deposition. Sputter deposition is an important thin film process

for a wide range of material systems and yet lacks the clear benefit of in situ diffraction that PLD

and MBE take advantage of through RHEED.

A novel thin film deposition system was designed and constructed to incorporate PLD,

sputtering, and RHEED into one chamber body. This design allowed for an easy switch between

PLD and sputter deposition all while maintaining a diffraction pattern. Using these growth

processes in conjunction with one another allowed for the deposition of complex epitaxial

heterostructures for superconducting quantum bits that included oxides and refractory metals,

which are extremely challenging to deposit via PLD and easily grown via sputtering. The

microwave loss properties of these shunt capacitors are extremely sensitive to the interface states

and as such RHEED was an invaluable tool for monitoring the interfaces in real time throughout

the deposition.

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While we were able to monitor the surface in real time during the multilayer deposition,

we were not able to observe intensity oscillations during the sputter deposition indicative of layer

by layer growth. To observe this powerful phenomenon during sputtering which is the most

commonly desired feature of real time RHEED analysis, a new material system was sought that

was known to grow layer by layer during PLD and also exhibited a strong thickness dependence.

To accomplish this, LaAlO3 growth was optimized separate from the RHEED chamber and was

shown to have electric properties similar to PLD grown films. This was the first time that a

conducting 2DEG was formed at the LAO/STO interface by sputtering and showed a sharp

transition to the conducting state after 4 unit cells of growth. This laid the ground work for

future studies of layer by layer controlled growth with in situ RHEED.

Once the conducting interface was formed, the LAO growth was moved to the hybrid

chamber where real time RHEED analysis was possible. Despite the scattering due to high

pressures of argon and also the strong deflection of the beam due to the magnetron sputter

source, clear RHEED oscillations were observed during the 90° off-axis LAO growth. This was

the first demonstration of layer by layer controlled growth by RHEED during sputter deposition

that we are aware of. To further improve on this process and also demonstrate the robustness of

the technique, LSMO and SRO films were deposited and both showed clear, large amplitude

intensity oscillations out to greater than 20 unit cells exhibiting the ability to have precise unit

cell control during sputter deposition. Interestingly, this was also the first time that extended

RHEED oscillations (>10) had been observed during SRO growth which is commonly known to

grow in the step flow regime during sputtering. This is something that may have interesting

implications and deserves further investigation in the future.

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As acknowledged here and in previous works, the magnetic fields produced by the sputter

sources can have a large negative impact on the electron beam used for RHEED analysis.

Depending on the growth setup, the magnetron sputter sources can bend the beam so far that it is

no longer detectable. This was what we observed during the 90° off-axis deposition.

Fortunately, the deposition chamber had sufficient degrees of freedom to accommodate this

deflection and obtain the RHEED pattern back on the phosphor screen. However, this is not a

permanent solution and not a plausible one for all growth geometries. As such, finite element

modeling was used to better predict the deflection of the electron beam and also used as a tool to

design ways of avoiding the massive deflections.

Through the finite element modeling, we were able to predict the deflection we observed

in experiment and proceed to identify a setup that significantly mitigates the deflection and

restricts it to a uniaxial and predictable bending. This is accomplished by adding a second

sputter source to the chamber that is facing the original source. Most importantly, the magnets in

the second sputter source need to be rearranged such that they are oppositely poled to the first

source, antisymmetric, which results in horizontal field lines near the sample surface and a

strictly vertical deflection of the beam. This was then expanded to be applicable to all two gun

setups that have some degree of non on-axis nature to them as long as the sputter sources have

antisymmetric magnetic polarities. Furthermore, for on-axis deposition, a facing sputter source

behind the heater with symmetric polarity is needed to completely reduce the Lorentz force that

the electron beam experiences to zero. In this way, a large portion of sputter geometries are

accounted for and have a clear and straightforward path to mitigating or erasing the deflection of

the RHEED beam due to the magnetron sputter guns.

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This finding has vast implications, not only for creating a new ability to control film

thicknesses at the unit cell level in real time during sputter deposition, but also for potentially

getting even closer to Feynman’s prediction. Certain areas of material space may be

significantly easier to access through sputter deposition (nitrides, high oxidation states, etc.)

when compared to PLD, MBE, or other techniques. However, these areas may have gone

unexplored thus far due to the lack of necessary in situ tools for fine level control. New types of

superlattices may be possible, with the ability to have unit cell control of the SRO growth for

example, that were not abundantly feasible previously. One area of particular interest may be

moving towards true reactive sputtering with the incorporation of pure metal targets. The

alkaline earth metal oxides, for instance, are typically unstable in air, and as such can be very

difficult to use as target materials. Whereas, strontium metal for instance, can be used as a target

material and through reactive sputtering can form strontium oxide films. Combine this with a

ruthenium target and RHEED, and alternating deposition similar oxide MBE may be possible

where fine control over the layers and stoichiometry is possible. This area has largely been

unexplored. Additionally, higher order phases can be accessed in this way that are metastable

otherwise. This is only one specific example but the options are vast given the uncharted space

that exists due to the lack of a proper real time tool during sputtering, until now.