Ramanathan Sh. (Ed.) Thin Film Metal-Oxides: Fundamentals and Applications in Electronics and Energy

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Contents

1 In Situ Synchrotron Characterization of Complex Oxide

Heterostructures .............................................................. 1

Tim T. Fister and Dillon D. Fong

2 Metal-Insulator Transition in Thin Film Vanadium

Dioxide ......................................................................... 51

Dmitry Ruzmetov and Shriram Ramanathan

3 Novel Magnetic Oxide Thin Films .......................................... 95

Jiwei Lu, Kevin G. West, and Stuart A. Wolf

4 Bipolar Resistive Switching in Oxides for Memory

Applications ...................................................................131

Rainer Bruchhaus and Rainer Waser

5 Complex Oxide Schottky Junctions ........................................169

Yasuyuki Hikita and Harold Y. Hwang

6 Theory of Ferroelectricity and Size Effects in Thin Films ...............205

Umesh V. Waghmare

7High-T

Superconducting Thin- and Thick-Film–Based

Coated Conductors for Energy Applications ..............................233

C. Cantoni and A. Goyal

8 Mesostructured Thin Film Oxides ..........................................255

Galen D. Stucky and Michael H. Bartl

9 Applications of Thin Film Oxides in Catalysis ............................281

Su Ying Quek and Efthimios Kaxiras

xiii

xiv Contents

10 Design of Heterogeneous Catalysts and the Application

to the Oxygen Reduction Reaction..........................................303

Timothy P. Holme, Hong Huang, and Fritz B. Prinz

Index .................................................................................329

Contributors

Michael H. Bartl Department of Chemistry and Department of Physics,

University of Utah, Salt Lake City, UT, USA

Rainer Bruchhaus Forschungszentrum Juelich, Institute of Solid State Research,

Juelich, Germany

Claudia Cantoni Oak Ridge National Laboratory, Oak Ridge, TN, USA

Tim T. Fister Materials Science Division, Argonne National Laboratory, Argonne,

IL, USA

Dillon D. Fong Materials Science Division, Argonne National Laboratory,

Argonne, IL, USA

Amit Goyal Oak Ridge National Laboratory, Oak Ridge, TN, USA

Yasuyuki Hikita Department of Advanced Materials Science, Graduate School

of Frontier Sciences, University of Tokyo, Tokyo, Japan

Timothy P. Holme Mechanical Engineering Department, Stanford University,

Stanford, CA, USA

Hong Huang Department of Mechanical and Materials Engineering, Wright State

University, Dayton, OH, USA

Harold Y. Hwang Department of Advanced Materials Science, Graduate School

of Frontier Sciences, University of Tokyo, Tokyo, Japan

Efthimios Kaxiras Department of Physics, School of Engineering and Applied

Sciences, Harvard University, Cambridge, MA, USA

Jiwei Lu Department of Materials Science and Engineering, University

of Virginia, Charlottesville, VA, USA

Fritz B. Prinz Mechanical Engineering Department, Stanford University,

Stanford, CA, USA

Su Ying Quek School of Engineering and Applied Sciences, Harvard University,

Cambridge, MA, USA

xvi Contributors

Shriram Ramanathan School of Engineering and Applied Sciences, Harvard

University, Cambridge, MA, USA

Dmitry Ruzmetov School of Engineering and Applied Sciences, Harvard

University, Cambridge, MA, USA

Galen D. Stucky Department of Chemistry and Biochemistry and Materials,

University of California, Santa Barbara, Santa Barbara, CA, USA

Umesh V. Waghmare Theoretical Sciences Unit, Jawaharlal Nehru Centre

for Advanced Scientiﬁc Research, Jakkur, Bangalore, India

Rainer Waser Forschungszentrum Juelich, Institute of Solid State Research,

Juelich, Germany

Kevin G. West Department of Materials Science and Engineering, University

of Virginia, Charlottesville, VA, USA

Stuart A. Wolf Department of Materials Science and Engineering, University

of Virginia, Charlottesville, VA, USA

Chapter 1

In Situ Synchrotron Characterization

of Complex Oxide Heterostructures

Tim T. Fister and Dillon D. Fong

Abstract This chapter surveys the high temperature and oxygen partial pressure

behavior of complex oxide heterostructures as determined by in situ synchrotron

X-ray methods. We consider both growth and post-growth behavior, emphasizing

the observation of structural and interfacial defects relevant to the size-dependent

properties seen in these systems.

1.1 Introduction

Complex oxides display a remarkable range of materials properties: some are

insulating and ferroelectric while others demonstrate colossal magnetoresistance

or high-temperature superconductivity. In these multifunctional materials, charge,

spin, and orbital degrees of freedom couple with dynamical lattice effects, present-

ing a rich playground for both fundamental and applied research. Excitement in

this ﬁeld has continued to build following the recent discoveries of novel phenom-

ena exhibited by atomically tailored oxide heterostructures [1]. Symmetry breaking

and charge transfer across an oxide interface can lead to emergent properties – e.g.,

superconducting interfaces between insulating oxides [2] – while epitaxial strain al-

lows for the production of new structural phases or electronic ground states [3, 4].

Poor control over oxide synthesis, however, can result in irreproducibility and con-

tribute to signiﬁcant discrepancies in the scientiﬁc literature.

As many of these novel properties are strongly dependent on stoichiometry

and ﬁnite size effects, strict control over ﬁlm composition and thickness is essen-

tial. Common growth techniques include molecular beam epitaxy (MBE), pulsed

laser deposition (PLD), and metalorganic chemical vapor deposition (MOCVD).

With such deposition methods, it is possible to achieve reasonable stoichiometric

and excellent thickness control. Growth takes place far from equilibrium, however,

meaning that the kinetics of deposition and oxygen incorporation can strongly in-

T.T. Fister (



) and D.D. Fong (



)

Argonne National Laboratory, Materials Science Division, Argonne, IL, USA

e-mail: ﬁster,fong@anl.gov

S. Ramanathan (ed.), Thin Film Metal-Oxides: Fundamentals and Applications

in Electronics and Energy, DOI 10.1007/978-1-4419-0664-9

 Springer Science+Business Media, LLC 2010

2 T.T. Fister and D.D. Fong

ﬂuence the resulting structure and electronic properties. Small variations in laser

ﬂuence during PLD of SrTiO

, for example, can result in cation nonstoichiometry,

leading to signiﬁcant changes in its lattice constant and many order-of-magnitude

changes in its conductivity [5]. These strong structure–composition–property rela-

tionships are typical for many of the complex oxides, underscoring the need for in

situ structural and chemical probes during growth and post-growth processing.

In this chapter, we discuss in situ synchrotron X-ray scattering and spectroscopy

and their utility in the study of complex oxide heterostructures. X-rays, unlike

electron probes, interact weakly with samples. Scattering is therefore kinematic,

facilitating analysis, and the large attenuation length of hard X-rays permits studies

in the high temperature .T / and oxygen partial pressure .P

/ environments typical

for oxide synthesis and processing. Synchrotron X-rays are highly brilliant, tunable

in energy, polarized, and sent in ultrafast pulses, enabling a wide array of scatter-

ing and spectroscopic techniques [6]. Relevant examples include resonant studies of

charge, spin, and orbital ordering [7], correlation spectroscopy [8,9], and investiga-

tions of domain dynamics with nanosecond resolution [10]. Furthermore, with the

manifold advances in X-ray optics, detectors, and analytical tools [11–13], 3D real-

space imaging is becoming more commonplace, affording the ability to compare

ensemble-averaged information with real-space images of structural or chemical

properties mapped with atomic-scale resolution. This will be paramount in gaining

insight into how emergent properties arise from nanostructures or nanoscale defects.

Our present focus is on recent studies employing synchrotron methods for the ex-

amination of complex oxide heterostructures during deposition or high temperature

and pressure processing. Given the recent spate of activity in perovskite systems,

we further restrict ourselves to a discussion of oxides with this particular crystal

structure.

The text is organized as follows. In Sect. 1.2, we provide background on the

perovskite structure and X-ray scattering/spectroscopy. Studies on the synthesis of

complex oxide thin ﬁlms are presented in Sect. 1.3, focusing primarily on the growth

of perovskite ﬁlms on SrTiO

(001) substrates by PLD and MOCVD. Section 1.4

concerns the interface and through-thickness structure of oxide ﬁlms as determined

by model ﬁtting and phase-retrieval techniques. The formation of perovskite do-

mains as detected by diffuse scattering and spectroscopic studies on manganites and

titanates are also discussed. We conclude with a few words on the future impact of

in situ synchrotron studies on complex oxide heterostructures.

1.2 Background

1.2.1 Perovskites

Complex oxides exhibit strong structure–composition–propertyrelationships. Thus,

point, line, and planar defects are expected to strongly affect properties, particu-

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 3

Fig. 1.1 Four depictions of the perovskite structure as 2  2  2 pseudocubic cells. (a) Pm

(e.g., SrTiO

). (b) Pbnm (e.g., SrRuO

). (c) R

3c (e.g., La

0:7

0:3

MnO

). (d) P4mm (e.g., fer-

roelectric PbTiO

). The orthorhombic and rhombohedral unit cells are shown in (b)and(c),

respectively

larly in heterostructures where layer thicknesses are on the nanometer scale. With

regard to the perovskite ABO

, where A is an (alkaline) rare earth and B is a transi-

tion metal, rare earth, or group III metal, much of the interesting phenomena stems

from strong electron interactions between the partially ﬁlled d- or f-shells of the

A- or B-site cations as mediated by the surrounding oxygens. The ideal perovskite

structure is shown in Fig. 1.1a, which has a tolerance factor

 D

C r

(1.1)

equal to 1 (where r

,andr

are the ionic radii of the A-, B-, and O-site ions).

The B-site cations are centralized in oxygen octahedra (6-coordinated), while the

A-site cations are surrounded by 12 oxygen anions. As the size of the A-site cation

4 T.T. Fister and D.D. Fong

shrinks, the octahedra tilt in an effort to minimize the A-site coordination volume,

asshowninFig.1.1b. Similar tilting occurs when the size of A-site cation grows

(Fig. 1.1c), and a useful rule of thumb regarding their effect on the Bravais lattice

is [14]

0:75    0:89

„

ƒ‚ …

orthorhombic

0:89    1:00

„

ƒ‚ …

cubic

1:00    1:05

„

ƒ‚ …

rhombohedral

; (1.2)

although both orthorhombic and rhombohedral perovskites are often referenced to

pseudocubic axes, as in Fig. 1.1. (A much more rigorous method of predicting per-

ovskite crystal structures has been presented by Lufaso and Woodward [15].) For

reference, the pseudocubic lattice parameters for orthorhombic (Pbnm) perovskites

are given by

p;1

ortho

C b

ortho

(1.3)

and

p;2

D c

ortho

=2 (1.4)

while that for rhombohedral perovskites is given by

rho

=2 (1.5)

or if in hexagonal coordinates

hex

: (1.6)

As discussed by Glazer [16], perovskites can be modeled by following three modi-

ﬁcations to the idealized structure: (i) tilting of the octahedra, (ii) displacements of

the cations, and (iii) distortions of the octahedra, although (iii) is often associated

with (ii). An example of cation displacements is shown in Fig. 1.1d, depicting fer-

roelectric displacements along the [001] direction. Each structural perturbation is

typically accompanied by an evolution in the crystal ﬁeld of the B-site ion, produc-

ing strong coupling between electronic properties and crystal structure, the latter of

which can be manipulated by epitaxial strain in oxide heterostructures. Electronic

or ionic carrier concentrations are typically tuned through aliovalent substitution on

the A- or B-sites requiring ionic or electronic compensation (e.g., oxygen vacancies

or holes) [17]. Similar neutralizing mechanisms take place in oxidizing or reducing

environments, and the equilibrium defect concentrations can be plotted as a func-

tion of P

if the rate constants are known [18]. Deposited oxide heterostructures at

room temperature, however, are in partially frozen-in states [19], and other methods

for estimating point defect concentrations are required [20–23].

Additional “knobs” for tuning material properties become available at perovskite

heterointerfaces. Aside from novel strain- or size-stabilized phases, local electric

ﬁelds induced by band alignment [24, 25], polar interfaces [26], and accumulated

charged defects [27] lead to a space-charge distribution often difﬁcult to predict

or control. It is this distribution, however, that can give rise to unusual electrical

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 5

effects at the interface like quasi-2D electron gas behavior [28] or colossal ionic

conductivity [29]. Researchers have therefore emphasized the deposition of epitaxial

oxide layers, minimizing the number of misﬁt dislocations at the interface to better

address the impact of other, less well-controlled defects on interfacial properties. It

is here that synchrotron X-ray scattering and spectroscopy illustrate their advantage

as not only an in situ growth monitor, but also an extremely powerful characteriza-

tion tool for non-destructively probing the heterointerface structure and chemistry

with atomic-scale resolution.

1.2.2 Scattering

A variety of X-ray scattering techniques are useful for the study of epitaxial

oxide heterostructures including X-ray reﬂectivity [30, 31], surface X-ray diffrac-

tion (SXRD) [32–35], and X-ray microscopy [36, 37]. Our emphasis is on SXRD,

but as we have previously considered in situ synchrotron X-ray studies of ferroelec-

tric materials [38], only a brief review is presented here.

Once corrected for instrumental and geometrical effects [39], the scattered in-

tensity is approximately equal to the Fourier transform of the sample’s electron

density distribution times its complex conjugate. For partially coherent scattering

probes, the total intensity is equal to the sum of intensities over all coherently il-

luminated regions. Navigation in reciprocal space is accomplished by rotating the

sample (and hence reciprocal space) in coordination with the detector such that the

feature of interest lies on the Ewald sphere at q D k

 k

(Fig. 1.2). Since the ra-

dius of the Ewald sphere is inversely proportional to the X-ray wavelength, smaller

Fig. 1.2 Experimental geometry for SXRD measurements. (a) The geometry in real space, il-

lustrating the incident k

and exiting k

wavevectors and the diffractometer angles ; ı,and!.

(b) The geometry in reciprocal space, demonstrating the measurement of anti-Bragg intensity along

a CTR. The scattering vector q lies on the Ewald sphere. Rotation about the sample normal by !

allows the feature of interest to meet with the scattering vector

6 T.T. Fister and D.D. Fong

wavelengths permit sampling of a larger volume in reciprocal space and therefore

improved spatial resolution. During SXRD, the incidence angle, ˛, is ﬁxed and kept

near the critical angle for total external reﬂection, ˛

, to optimize signal from the

surface region.

SXRD relies on the measurement of crystal truncation rods (CTRs), features in

reciprocal space that stem from the abrupt truncation of the crystal by the surface.

If the interfaces are sufﬁciently smooth, a vast amount of information regarding the

structure of the ﬁlm may be retrieved by SXRD. Here we present a simple way

of calculating the CTR structure factor for a coherently strained ﬁlm on a cubic

substrate.

The structure factor for a single unit cell can be written as

unitcell

.q/ D

j D1

.q/e

B

.q=4/

iqr

; (1.7)

where f

.q/ is the q-dependent form factor, B

is the isotropic Debye–Waller factor

for atom j ,andN

is the number of atoms in the unit cell. Tabulated values of

.q/ are provided in [40], and the anomalous dispersion corrections are available

online [41].

The scattering factor for a single column of unit cells along the out-of-plane

direction can be determined by summing the product of the appropriate unit cell

structure factor and its phase factor over every nth layer along the z-direction:

CTR

.q/ D

nD0

unitcell

.q/e

.iq

z;n

=

; (1.8)

where r

z;n

is the distance to the nth layer and 

accounts for absorption [42]. It is

often useful to distinguish between scattering from a ﬁlm N

ﬁlm

unit cells thick

ﬁlm

.q/ D F

unitcell

ﬁlm

.q/

1  e

iN

ﬁlm

N

ﬁlm

c=

ﬁlm

1  e

iq

c=

ﬁlm

(1.9)

and scattering from a semi-inﬁnite substrate

sub

.q/ D F

unitcell

sub

.q/

1  e

iq

a=

sub

; (1.10)

where the cubic substrate has a lattice constant a, and the in-plane and out-of-plane

lattice constants for the ﬁlm are a and c, respectively. The interlayer spacing be-

tween the ﬁlm and substrate, , gives rise to an additional phase factor such that the

total scattering factor along a CTR is given by

CTR

.q/ D F

ﬁlm

.q/ C e

iq

..N

ﬁlm

1/cC/

..N

ﬁlm

1/cC/=

ﬁlm

sub

.q/: (1.11)