University of Southern California Dissertations and Theses

Synthesis and surface modification of perovskite-based nanocrystals for use in high energy density nanocomposites

(USC Thesis Other)

Synthesis and surface modification of perovskite-based nanocrystals for use in high energy density nanocomposites

PDF

Download a page range

Download transcript

Copy asset link

Request this asset

Transcript (if available)

Content

SYNTHESIS AND SURFACE MODIFICATION OF PEROVSKITE-BASED
NANOCRYSTALS FOR USE IN HIGH ENERGY DENSITY NANOCOMPOSITES

by

Christopher W. Beier

________________________________________________________________________

A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(CHEMISTRY)

August 2013

Copyright 2013 Christopher W. Beier
ii

Acknowledgements
I thank my research advisor, Professor Richard Brutchey for his infallible support,
mentorship, and patience. Without his desire to challenge me both as a scientist and as a
person, my full potential would not have been realized. For this I am eternally grateful. I
would like to thank all the members of the Brutchey group, past and present, for their
friendship and fellowship over the years.
I thank Professors Mark Thompson, Andrea Hodge, G. K. Surya Prakash, Steven
Bradforth, and Hanna Reisler for serving as committee members and mentors as I worked
towards earning my doctorate. I would also like to acknowledge the Chemistry
Department and the David and Dana Dornsife College of Letters, Arts and Sciences.
I would also like to thank my wonderful wife, Alejandra, for always being there for me
no matter how hard things got or how unobtainable my goals seemed. She is the love of
my life, and I could not have made it without her support, love, and uncanny ability to
always make me see the best in myself. Finally, I would like to thank my mother, father,
and brother for being supportive of the ups and downs of this arduous process.

iii

Table of Contents
Acknowledgements ii
List of Figures vi

List of Schemes x

Abstract xi

Chapter 1. The Effects of Barium Titanate Nanocrystals in
Nanocomposite-Based Dielectrics

1
1.1. Introduction 1
1.2. Nanocomposite Preparation via Direct Polymer Addition 3
1.2.1. Direct Addition of Unmodified BaTiO
3
4
1.2.2. Direct Addition of Modified BaTiO
3
9
1.3.Nanocomposite Preparation via In Situ Polymerization 15
1.3.1. In Situ Polymerization in the Presence of
Unmodified BaTiO
3

15
1.3.2. In Situ Polymerization in the Presence of
Modified BaTiO
3

18
1.4. Conclusions and Perspective 25
1.5. References 26

Chaper 2. Low-Temperature Synthesis of Solid-Solution Ba
x
Sr
1-x
TiO
3

Nanocrystals

31
2.1. Introduction 31
2.2. Experimental Section 34
2.2.1. General Procedures 34
2.2.2. Synthesis of Ba
x
Sr
1-x
TiO
3
(0 ≤ x ≤ 1)
Nanocrystals

34
2.2.3. Material Characterization 35
2.2.4. Dielectric Characterization 35
2.3. Results and Discussion 36
2.3.1. Nanocrystal Synthesis and Characterization 36
2.3.2. Dielectric Properties 40
2.4. Conclusions 45
2.5. References 46

Chapter 3. Effect of Surface Modification on the Dielectric Properties
of BaTiO
3
Nanocrystals

49
3.1. Introduction 49
iv

3.2. Experimental Section 51
3.2.1. Synthesis of BaTiO
3
Nanocrystals 52
3.2.2. Surface Functionalization with n-
Hexylphosphonic Acid

52
3.2.3. Material Characterization 52
3.2.4. Dielectric Characterization 53
3.3. Results and Discussion 54
3.3.1. Surface Modification of BaTiO
3
Nanocrystals 54
3.3.2. Dielectric Studies 59
3.4. Conclusions 63
3.5. References 64

Chapter 4. Improved Breakdown Strength and Energy Density in
Thin Film Polyimide Nanocomposites with Small Barium
Strontium Titanate Nanocrystal Fillers

67
4.1. Introduction 67
4.2. Experimental Section 70
4.2.1. General Procedures 70
4.2.2. Vapor Diffusion Sol-Gel Synthesis of
Ba
0.7
Sr
0.3
TiO
3
(BST) Nanocrystals

70
4.2.3. Material Characterization 71
4.2.4. Device Preparation 73
4.2.5. Device Characterization 74
4.3. Results and Discussion 74
4.3.1. Synthesis, Characterization, and Dielectric
Properties of BST

74
4.3.2. PMDA-BAPB/BST Nanocomposite Fabrication
and Characterization

77
4.3.3. Dielectric Characterization 81
4.4. Conclusions 85
4.5. References 86

Chapter 5. Size- and Ligand-Dependent Dielectric Properties of
pDCPD/BaTiO
3
Nanocomposites

90
5.1. Introduction 90
5.2. Experimental Section 93
5.2.1. Surface Functionalization of BaTiO
3
with 10-
Undeceneoic Acid

93
5.2.2. Material Characterization 94
5.2.3. Device Preparation 95
5.2.4. Device Characterization 96
5.3. Results and Discussion 97

v

5.3.1. Surface Modification and Characterization of
BaTiO
3
Nanocrystals

97
5.3.2. pDCPD/BaTiO
3
Nanocomposite Fabrication and
Characterization

99
5.3.3. Dielectric Characterization 102
5.4. Conclusions 106
5.5. References 106

Bibliography 110

vi

List of Figures
Figure 1.1: Frequency dependence of the (a) dielectric constant
and (b) dielectric loss of PVDF/BaTiO3
nanocomposites at 5 vol% loading with various
particle sizes ranging from 25 to 500 nm.

6

Figure 1.2: TEM images of PMMA/BaTiO
3
displaying shell
thicknesses of (a) 7.5 nm, (b) 10 nm, (c) 15 nm, and
(d) 17 nm.

21

Figure 2.1: (a) Powder X-ray diffraction pattern of the cubic
SrTiO
3
nanocrystals. (b) Low-resolution TEM image
of the 5.9 nm SrTiO
3
and high-resolution TEM image
of a single SrTiO
3
nanocrystal shown at the inset.

36

Figure 2.2: (a) Powder X-ray diffraction pattern of the (110)
reflection for the Ba
x
Sr
1-x
TiO
3
nanocrystals (0 ≤ x ≤ 1)
referenced to the (111) reflection of an internal silicon
standard (*). (b) Linear dependence of the lattice
parameter as a function of composition for the Ba
x
Sr
1-
x
TiO
3
nanocrystals.

38

Figure 2.3: Low-resolution TEM images of (a) Ba
0.14
Sr
0.86
TiO
3
,
(b) Ba
0.27
Sr
0.73
TiO
3
, (c) Ba
0.33
Sr
0.67
TiO
3
, (d)
Ba
0.43
Sr
0.57
TiO
3
, (e) Ba
0.48
Sr
0.52
TiO
3
, (f)
Ba
0.63
Sr
0.37
TiO
3
, (g) Ba
0.69
Sr
0.31
TiO
3
, and (h)
Ba
0.77
Sr
0.23
TiO
3
nanocrystals.

38

Figure 2.4: (a) Selected area electron diffraction pattern for an
ensemble of Ba
0.69
Sr
0.31
TiO
3
nanocrystals and (b)
intensity line profile for lattice fringes of a single
Ba
0.69
Sr
0.31
TiO
3
nanocrystal (inset).

39

Figure 2.5: Nitrogen adsorption (red)-desorption (black)
isotherms for fractured pellets of (a) SrTiO
3
, (b)
Ba
0.69
Sr
0.31
TiO
3
, and (c) BaTiO
3
nanocrystals.

42

vii

Figure 2.6: Compositional dependence of the dielectric constant
of Ba
x
Sr
1-x
TiO
3
nanocrystals measured at 1 kHz.
Measurements were taken at 20 °C.

43

Figure 2.7: (a) Dielectric constant of Ba
0.69
Sr
0.31
TiO
3
, BaTiO
3
,
and SrTiO
3
nanocrystals as a function of frequency.
(b) Dielectric loss of Ba
0.69
Sr
0.31
TiO
3
, BaTiO
3
, and
SrTiO
3
.

44

Figure 3.1: (a) Low-resolution TEM image of 6 nm HPA-BaTiO
3

nanocrystals and high-resolution TEM image of a
single HPA-BaTiO
3
nanocrystal as inset. (b) X-ray
diffraction pattern of cubic HPA-BaTiO
3
.

54

Figure 3.2: TGA curves of unmodified and HPA-BaTiO
3

nanocrystals.

55

Figure 3.3: Nitrogen adsorption-desorption isotherm for (a)
unmodified and (b) HPA-BaTiO
3
nanocrystals.

56

Figure 3.4: High resolution XPS spectrum of the P 2p region of
HPA-BaTiO
3
.

56

Figure 3.5: FT-IR spectra of HPA, HPA-BaTiO
3
, and BaTiO
3

nanocrystals.

57

Figure 3.6:
31
P MAS NMR spectra of HPA and HPA-BaTiO
3
. 58

Figure 3.7: (a) Dielectric constant of unmodified BaTiO
3
and
HPA-BaTiO
3
as a function of frequency. (b)
Dielectric loss of unmodified BaTiO
3
and HPA-
BaTiO
3
as a function of frequency. Measurements
were taken at 25 °C.

60

Figure 3.8: (a) Dielectric constant of unmodified BaTiO
3
and
HPA-BaTiO
3
as a function of temperature. (b)
Dielectric loss of unmodified BaTiO
3
and HPA-
BaTiO
3
as a function of temperature. Measurements
were taken at a frequency of 2 kHz.

61

Figure 4.1: Schematic of high voltage testing station used for all
breakdown measurements.

75

viii

Figure 4.2: X-ray diffraction patterns of BST and PMDA-
BAPB/BST nanocomposite thin films.

76

Figure 4.3: (a) Low resolution TEM image of BST nanocrystals,
and (b,c) high-resolution TEM images showing the
{110} and {100} family of lattice planes.

76

Figure 4.4: Digital photograph of all PMDA-BAPB/BST
nanocomposites.

77

Figure 4.5: SEM micrographs with Ti elemental mapping (shown
as inset) for PMDA-BAPB/BST nanocomposites at 5
(a), 10 (b), 13 (c), 15 (d), and 18 (e) vol% BST
loading.

78

Figure 4.6: TGA curves in air (a) and in nitrogen (b) for PMDA-
BAPB/BST nanocomposites at all loadings.

79

Figure 4.7: FT-IR spectrum of the neat PMDA-BAPB polyimide. 80

Figure 4.8: Raman spectra of PMDA-BAPB/BST
nanocomposites, offset for clarity. The neat PMDA-
BAPB film was processed and prepared in an
analogous way to the nanocomposites but without
addition of BST.

80

Figure 4.9: (a) The Weibull breakdown at 63.2 % probability of
failure (red, left axis) and relative permittivity
compared to Bruggeman’s effective medium model
(blue, right axis) as a function of BST loading. (b)
Two-parameter Weibull plots of PMDA-BAPB/BST
nanocomposites where dashed line represents Weibull
breakdown at 63.2 % probability of failure.
Measurements were taken at 25 °C.

82

Figure 4.10: (a) Relative permittivity of PMDA-BAPB/BST
devices as a function of frequency. (b) Dielectric loss
of PMDA-BAPB/BST devices as a function of
frequency. Measurements were taken at 25 °C.

85

Figure 5.1: Powder XRD pattern of (a) BT50 powders and
pDCPD/BT50 nanocomposite films and (b) BT100
powders and pDCPD/BT100 nanocomposite films.

97
ix

Figure 5.2: FT-IR spectra of BT50, BT100, mBT50, and mBT100
nanocrystals.

98

Figure 5.3: Digital photograph of pDCPD/BT100 demonstrating
film flexibility.

99

Figure 5.4: SEM micrographs of (a) pDCPD/BT50, (b)
pDCPD/BT100, (c) pDCPD/mBT50, and (d)
pDCPD/mBT100.

100

Figure 5.5: TGA curves in nitrogen for all (a) pDCPD/BT50 and
(b) pDCPD/BT100 nanocomposites at 5 vol% loading,
and all (c) BT50 and (d) BT100 nanocrystals.

101

Figure 5.6: (a) Relative permittivity and dielectric loss of pDCPD,
pDCPD/BT50, and pDCPD/mBT50 devices as a
function of frequency. (b) Relative permittivity and
dielectric loss of pDCPD, pDCPD/BT100, and
pDCPD/mBT100 devices as a function of frequency.
Measurements were taken at 25°C.

103

Figure 5.7: Two-parameter Weibull plots of pDCPD/BaTiO
3

nanocomposites where the dashed line represents
Weibull breakdown at 63.2 % probability of failure.
Measurements were taken at 25 °C.

105

x

List of Schemes
Scheme 1.1: Proposed hydrogen bonding interaction with PVDF in the
interfacial zone.

5

Scheme 1.2: Schematic illustration of the formation of BaTiO
3

nanoparticles with the polymer shell.

10

Scheme 1.3: Proposed interaction of PAA oligomers with the surface of
BaTiO
3
.

16

Scheme 1.4: Schematic diagram illustrating the process of ATRP from
the surface of BaTiO
3
.

19

Scheme 1.5: Schematic representation of BaTiO
3
surface modification,
grafting, and curing in the formation of bisphenol A based
nanocomposite.

23

xi

Abstract
Perovskite-based oxides like BaTiO
3
, SrTiO
3
, and Ba
x
Sr
1-x
TiO
3
are extremely useful
materials which possess fundamental properties such as high dielectric constants, high
energy densities, and low loss tangents. In addition, perovskites display size,
composition, and synthesis dependent dielectric properties. As such, they have found
themselves at the forefront of modern capacitive and charge storage applications. While
promising, such materials are typically prepared at high temperatures (>500 °C), are
extremely brittle, and cannot be easily incorporated into flexible devices. In order to
overcome the limitations associated with the processing of high permittivity ceramics,
researchers have sought to use a nanocomposite approach, whereby small, well-defined
perovskite nanocrystals are integrated into a polymer matrix. In this approach, the
polymer provides processability and high breakdown strength (E
bd
), while the perovskite
filler delivers improved dielectric performance. With this in mind, we developed a low
temperature route to preparing gram-scale quantities of small (<15 nm), well-defined
BaTiO
3
, SrTiO
3
,

and Ba
x
Sr
1-x
TiO
3
nanocrystals and blended them into novel polymeric
systems. Through precise control of the nanocrystal composition and surface chemistry,
it was possible to better understand what factors govern the dielectric performance of the
nanocrystals. Compositionally, the dielectric constant of the nanocrystal can be increased
by more than an order of magnitude, while controlled surface modification enhances
dispersion and improves dielectric temperature and frequency stability. This information
was then used to prepare novel composite systems with enhanced performance.
1

Chapter 1. The Effects of Barium Titanate Nanocrystals in
Nanocomposite-Based Dielectrics

1.1 Introduction
Barium titanate (BaTiO
3
) was first synthesized between 1924 and 1926 by V. M.
Goldschmidt and has since become one of the most notable and functional members of
the perovskite family of materials.
1
Bulk BaTiO
3
exhibits ferroelectric behavior as a
result of a distortion of the ideal perovskite structure below 130 ˚C. In this non-
centrosymmetric, tetragonal phase (space group P4mm), the Ti
4+
cation is displaced from
the center of the unit cell relative to one of six O
2–
anions, which results in a net dipole
moment that can be oriented from one crystallographic direction to another via an applied
electric field.
2,3
Upon heating bulk BaTiO
3
to its Curie temperature (T
C
) at 130 ˚C, the
dipolar properties are lost and the material becomes paraelectric as a result of a reversible
phase transformation to the ideal centrosymmetric, cubic phase (space group Pm3m). In
addition to its ferroelectric properties, BaTiO
3
also possesses relative permittivities as
high as ε
r
= 10,000 and is piezoelectric and pyroelectric, all of which make it a very
desirable technological material for applications in capacitors, nonvolatile memory,
uncooled IR detection, and actuators.
4
Over the past several years there have been
several attempts to synthesize BaTiO
3
nanocrystals, which display a range of size-
dependent ferroelectric, dielectric, and structural characteristics.
5–12
It has been
theoretically predicted, and experimentally confirmed, that for nanoscale BaTiO
3
T
C
is
suppressed with decreasing particle size, and the paraelectric cubic phase becomes
thermodynamically preferred at lower temperatures due, in part, to increased interfacial
2

area and surface energy.
13–15
While this trend is consistent, the absolute dielectric
constant is highly dependent upon preparative methods. For instance, BaTiO
3

nanocrystals 50-70 nm in diameter can display relative permittivities ranging from ε
r
=
180 to 2500.
16–19
Additionally, the non-centrosymmetric tetragonal phase can be seen in
particles as small as 4 nm at room temperature
20
, despite studies which suggest the cubic
phase is preferred in grain sizes below 30 nm.
16,21

Lately, several attempts have been made to take advantage of the unique dielectric
properties of nanocrystalline BaTiO
3
by blending them into polymers, forming
nanocomposites.
22,23
The underlying motivation behind using these hybrid
nanocomposites is that the polymer component will contribute high breakdown strength,
solution processability, and flexibility while the nanoparticulate filler will provide a high
dielectric constant. The maximum energy density for a capacitor is given by U
v
= ½ ε
r
ε
0
E
BD
2
(J cm
–3
), where ɛ
r
is the relative permittivity of the nanocomposite, ε
0
is the
permittivity of free space, and E
BD
is the breakdown strength of the nanocomposite.
24

While there is enormous potential in using polymer/BaTiO
3
composites to fabricate the
next generation of high energy density capacitors, significant gains have yet to be
realized. The large interfacial area of nanocrystals leaves a considerable volume fraction
of polymer that is affected by the nanocrystal surface and displays properties different
from that of the bulk polymer. The region where the nanocrystals have an intimate
interaction with the polymer is called the interfacial zone.
25
Because of the interfacial
zone, adding only a few weight percent of nanocrystalline BaTiO
3
can cause atypical
changes in permittivity and breakdown strength, and it can change the chemical and
3

physical stability of the nanocomposite before significant improvements are realized.
23,26–
28

To date, there is no unifying theory which explains the role of the interfacial zone in
BaTiO
3
-polymer nanocomposites and current composites depend upon system-by-system
optimization utilizing a variety of physical (e.g., ball milling, ultrasonic processing, high
shear mixing) and chemical routes (e.g., surface grafting, functionalization, surface
pretreatment) to mitigate agglomeration, improve dispersion, and optimize device
performance.
29
Herein, recent advancements in polymer/BaTiO
3
nanocomposites will be
summarized in order to understand how nanoparticle size, processing conditions, surface
chemistry, and choice of polymer affect the composite dielectric properties. The
interfacial zone will also be studied in order to define its role in optimizing device
performance.
1.2 Nanocomposite Preparation via Direct Polymer Addition
The most straight forward route to preparing BaTiO
3
nanocomposites is to disperse the
nanocrystals directly into a solution containing a pre-synthesized polymer.
Nanocomposite films can be prepared by casting solutions of varying concentrations into
containers or coating them on substrates, with the film forming after solvent evaporation.
Ultrafine (<100 nm) BaTiO
3
nanoparticles display inherently high surface energies and
chemical activity compared to larger micron sized particles, and an understanding of their
interactions within the interfacial zone is critical in the decision of proper polymer,
solvent, and processing conditions.
29,30
A range of methods have been recently employed
4

that seek to take advantage of the chemically reactive surface of BaTiO
3
by either
working with BaTiO
3
in its natural state or by altering its surface to enhance or control
the interfacial interaction and improve dielectric properties. The effects of directly
mixing both modified and unmodified BaTiO
3
into a range of polymer systems will be
examined in the following sections with special attention paid to the interfacial effects on
the composite’s dielectric performance.
1.2.1 Direct Addition of Unmodified BaTiO
3

The most widely studied polymer for BaTiO
3
-based nanocomposites is poly(vinylidene
fluoride) (PVDF), a solution processable semicrystalline thermoplastic polymer.
31,32

PVDF and its copolymers are polarizable low melting ferroelectrics with dielectric
constants as high as ε
r
= 52, making them highly desirable candidates for dielectric
applications.
32
PVDF and its derivatives are soluble in polar aprotic solvents such as
N,N’-dimethylformamide (DMF), tetrahydrofuran (THF), and methyl ethyl ketone
(MEK). PVDF can crystallize in a variety of phases, but the most prevalent are the α-
phase and β-phase. The α-phase is thermodynamically preferred but is non-polar and has
low permittivities, while the β-phase is polar and displays much higher permittivities.
The phase can be controlled through processing conditions and thermal treatment of the
films.
32
In an attempt to take advantage of PVDF’s interesting dielectric properties,
Mendes et al. sought to study the size dependent effects of PVDF/BaTiO
3
composites
with BaTiO
3
nanocrystals that were 10 and 500 nm in diameter. To make the composite,
BaTiO
3
solutions in DMF were dispersed with sonication, mixed with a measured
5

amount of PVDF powder (amount dependent upon desired loading, 5 or 10 % BaTiO
3
by
weight), and cast into films.
33
Fourier transform infrared (FT-IR) analysis of the as-
prepared films revealed that the smaller 10 nm BaTiO
3
nanocrystals acted as a nucleation
site for preferred growth of the β-phase of PVDF, indicated by a decrease in the relative
ratio of adsorption bands at 764 and 840 cm
-1
corresponding to the α- and β-phases,
respectively. The 500 nm nanocrystals, on the other hand, still acted as a nucleation site
but had a much smaller concentration of the β-phase. Surface hydroxyl groups at the
surface of BaTiO
3
formed a strong O-H
…
F-C hydrogen bonding interaction at the
BaTiO
3
/PVDF interface, allowing for crystal growth and alignment, shown in Scheme
1.1. The enhanced interaction of the smaller 10 nm nanocrystals was independent of
concentration, and found to be directly proportional to interfacial area, which is
approximately 50 times larger than the 500 nm BaTiO
3
. Interestingly, the permittivity of
the BaTiO
3
/PVDF composites were highest, ε
r
= 75, for the 10 nm particles instead of the
500 nm particles at the same loading. Given that the permittivity of BaTiO
3
decreases as
the particle size decreases (vide supra), it is expected that the composite permittivity
should reduce as the particle size decreases. The higher permittivity of the composites

Scheme 1.1. Proposed hydrogen bonding interaction with PVDF in the interfacial zone.
6

made with smaller nanocrystals was proposed to be a result of the increased interfacial
interaction between BaTiO
3
and PVDF which improved space charge distribution.
33
As
the size of the nanocrystal decreased, the interface between the ceramic and the polymer
dominated the dielectric response, showing that the composite permittivity did not only
rely on the absolute permittivity of the matrix and filler. Work from Mao, Dang, and
Kobayashi also found an atypical dependence of the dielectric permittivity on particle
size for BaTiO
3
/PVDF composites, reaching a maximum permittivity with smaller
BaTiO
3
nanocrystals at low filler concentrations (10 vol% or less) (Figure 1.1).
34–36

Dang and Kobayashi additionally discovered that the atypical behavior in these systems
disappeared as the concentration of BaTiO
3
was raised, increasing the interparticle
contact and mitigating the interfacial effects of PVDF.
34

In an attempt to understand the role of the polymer, Pant and coworkers studied the

Figure 1.1. Frequency dependence of the (a) dielectric constant and (b) dielectric loss of
PVDF/BaTiO
3
nanocomposites at 5 vol% loading with various particle sizes ranging from 25 to
500 nm. Modified from Reference 35.
7

concentration dependent microwave frequency dielectric properties of BaTiO
3

nanocomposites using two polar polymers: polyaniline (PANI) and maleic resin.
37
The
high frequency permittivities of ferroelectric materials, like BaTiO
3
, depend largely on
dielectric relaxation and dispersion phenomena, grain boundaries, and particle size. It
was of interest in their study to determine if microwave dependent relaxation phenomena
in BaTiO
3
could be tuned through the composite approach. Structural analysis of the
composites displayed significant concentration dependence on the observed physical
properties of the polymer. FT-IR analysis showed that as the concentration of BaTiO
3

increased in PANI/BaTiO
3
composite system, N-H stretching modes consistent with
PANI began to dissipate. The reduced N-H intensity was determined to be caused by the
N-H groups forming a strong interaction at the surface of the nanoparticles. Additionally,
in the presence of BaTiO
3
, PANI became XRD amorphous. The maleic resin/BaTiO
3

composite, on the other hand, had a much weaker interaction and revealed no
recognizable chemical or structural changes in the presence of BaTiO
3
. Both systems
showed frequency independent permittivity, likely a result of an increase in the space
between individual BaTiO
3
grains in the composite that mitigated boundary induced high
frequency effects and stabilized the permittivity.
37
To further study the concentration-
induced dielectric effects, Pant and coworkers compared the relative permittivites against
a range of effective medium models which mathematically predict the dielectric response
in multi-component homogeneously dispersed systems.
38
It was found that the maleic
resin/BaTiO
3
system fit the asymmetric Bruggeman model
38
across the entire range,
8

whereas the BaTiO
3
-PANI system could not be fitted to any known model, which was
attributed to the interaction occurring at the interfacial zone.
37

Using a novel electrolytic deposition (ELD) technique, Kim and coworkers were able to
successfully prepare poly(methyl methacrylate) (PMMA)/BaTiO
3
composites at loadings
as high as 60 vol%.
39
By taking advantage of the hydroxyl coated surface of unmodified
BaTiO
3
, they were able to successfully attract homogeneous dispersions of BaTiO
3
to the
surface of anodic copper foil using an electric field of 50-100 V cm
–1
. By adding varying
amounts PMMA into the BaTiO
3
solution during deposition, it was possible to co-
precipitate both components onto the copper foil at a range of loadings. While it was not
determined if PMMA participated in a chemical reaction or simply physisorbed at the
surface of BaTiO
3
, it was found to be essential in creating an evenly dispersed non-
agglomerated film by forming a pseudo core/shell like structure around BaTiO
3
. The
core/shell structure isolated individual particles through either electrostatic or
electrosteric stabilization and aided in particle isolation while being deposited. The facile
approach of this method allowed for films to be created with loadings as high as 60 vol%
BaTiO
3
, thereby maximizing permittivity at ε
r
= 25.
39

More recently, Tang and coworkers prepared flexible and thermally stable BaTiO
3

nanocomposites in polyarylene ether nitrile (PEN) through a novel continuous ultrasonic
dispersion fabrication process at nanocrystal loadings up to 40 wt%.
40
As the
concentration of BaTiO
3
increased to 40 wt% in PEN, there was a 220% increase in the
relative permittivity, which could not be predicted by traditional effective medium
9

models (Bruggemann or Lichteneker’s law).
38
The breakdown strength consistently
decreased as more BaTiO
3
was incorporated into the system, but when combined with the
favorable permittivity data, increased the energy density by approximately 30% from
pure PEN. Further electrical characterization revealed that the composites were
frequency independent up to 200 kHz, and had volumetric resistivity values comparable
to PEN at all loadings. It is postulated that the uncharacteristic permittivity increase
arose due to BaTiO
3
dipole and interfacial polarization, which was not accounted for in
the aforementioned models.
40
Though not considered in this paper, the Jayasundere and
Smith model takes into account the dipolar effects of neighboring dielectric bodies and
has successfully predicted such effects in epoxy/BaTiO
3
composites (vide infra).
41

1.2.2 Direct Addition of Modified BaTiO
3

In an attempt to take advantage of the strong hydrogen bonding interaction of BaTiO
3
,
several studies were performed in which the surface of BaTiO
3
was “activated” with
hydrogen peroxide. By refluxing BaTiO
3
in concentrated hydrogen peroxide it is
possible to create a more densely packed layer of hydroxyl groups which are proposed to
enhance interfacial chemistry and improve dispersion over particles which did not receive
such treatment.
42,43
Zhou and coworkers studied BaTiO
3
activation with commercial
particles of less than 100 nm in diameter, using PVDF as the matrix material.
42

Hydrogen peroxide treated particles had consistently lower relative permittivity than the
unmodified BaTiO
3
at all loadings tested, but displayed better frequency stability due to
the increased interfacial hydrogen bonding at the activated nanocrystal interface that
10

restricted chain mobility and reduced composite polarization.
42
Further analysis of the
dielectric breakdown strength showed a marked enhancement of the activated over the
unmodified nanocrystal-containing composites. A similar result was found by
Almadhoun and coworkers, who also reported improved breakdown voltages in the
hydrogen peroxide treated PVDF/BaTiO
3
nanocomposites which they attributed to a
lower leakage current density than composites with unmodified BaTiO
3
.
43

In an analogous attempt to take advantage of the surface chemistry of BaTiO
3

nanocrystals, Jung and coworkers activated the nanocrystals under UV-ozone irradiation
then washed the particles with NaH, creating an anionic surface.
44
The activated BaTiO
3

was mixed with polystyrene-block-poly(styrene-co-vinylbenzylchloride) (PS-b-PSVBC),
a diblock copolymer, or poly(styrene-co-vinyl-benzylchloride) (r-PSVBC), a random

Scheme 1.2. Schematic illustration of the formation of BaTiO
3
nanoparticles with the polymer
shell. Modified from reference 44.
11

copolymer, and treated with triethylamine, forming a solution processable core/shell
heterostructure shown in Scheme 1.2. Solutions were spun cast into films, using neat
polystyrene to control the composite concentration. The diblock copolymer shell
outperformed the random copolymer system, displaying a higher permittivity (ε
r
= 45),
lower dielectric loss (tan δ = 0.06), and higher breakdown strength (222 V μm
–1
) at
loadings up to 75 wt% BaTiO
3
. The improved dielectric properties of this system were a
result of the insulating polymeric shell which effectively buried the charge at the
nanocrystal surface and reduced charge carrier mobility.
44

The inherently reactive surfaces of BaTiO
3
can also be organically modified by
carboxylates, silanes, or phosphonates in order to improve organic or aqueous solvent
dispersion, reduce particle aggregation, and/or to tune dielectric properties, further
expanding the utility of BaTiO
3
.
45,46
Perry and coworkers have recently functionalized
the surface of commercial 50 nm BaTiO
3
nanoparticles with pentafluorobenzyl
phosphoric acid (PFBPA) to prepare nanocomposites with poly(vinylidene fluoride-co-
hexafluorpropylene) (PVDF-HFP), a PVDF co-polymer. The PFBPA-BaTiO
3

nanocrystals were mixed with a controlled amount of PVDF-HFP and ball-milled for 14
days using DMF as a solvent and then cast into films.
22,47
The relative permittivity of the
composites were measured as a function of nanocrystal loading and were found to
increase continuously to ε
r
= 37 at 60 vol% BaTiO
3
and fell as the concentration was
increased further. The dielectric breakdown strength, on the other hand, remained above
300 V μm
–1
up to 10 vol%, and fell precipitously between 10 and 20 vol%, before
leveling off at about 225 V μm
–1
at higher loadings. By using a statistical packing model,
12

they were able to reasonably predict the observed trends in permittivity and breakdown
strength. The non-linear response of ε
r
was a result of the formation of air voids as the
PFBPA-BaTiO
3
content increased. At a critical value of 39.7 vol% BaTiO
3
, no
additional nanocrystals could be added without the removal of polymer, introducing
porosity. By accounting for localized field enhancements in this model, they were better
able to fit the experimental breakdown data, but could still not explain the precipitous
drop above 10 vol%.
22
Despite this drop, the maximum calculated energy density at 10
vol% BaTiO
3
was 8 J cm
–3
, about double that of the pure polymer.
24

Li and coworkers achieved a similar enhancement by preparing 60 nm BaTiO
3

nanocrystals via decomposition of ethylene diamine modified Ti(OiPr)
2
and Ba(OH)
2
in
an aqueous solution. The synthetic approach functionalized the 50-70 nm BaTiO
3

nanocrystal surface with ethylene diamine, rendering them soluble in a range of
solvents.
48
The ethylene diamine-coated BaTiO
3
nanocrystals were blended into two
PVDF copolymers, poly(vinylidene fluoride-co-chlorotrifluoro-ethylene) (P(VDF-
CTFE)) and poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene)
(P(VDF-TrFE-CTFE)), with permittivities of ε
r
= 12 and 42, respectively. They sought
to study the effects of polymer permittivity on the dielectric response of the composite
and found that the composite permittivity steadily increased for both systems reaching a
maximum of ε
r
= 24 and 50 at 20 vol% ethylenediamine-BaTiO
3
for the P(VDF-CTFE)
and P(VDF-TrFE-CTFE) copolymers, respectively. When comparing the composite
energy density for both systems, it was found that P(VDF-TrFE-CTFE)/BaTiO
3

improved 120%, whereas P(VDF-CTFE)/BaTiO
3
only produced an increase of about
13

97%. They believed that by matching the high permittivity of the nanocrystals (ε
r
= 180)
more closely to the polymer, they mitigated localized field enhancements and improved
composite polarizability.
48
These local field inhomogeneities occured as a result of the
drastically different permittivities between BaTiO
3
and the polymer matrix, which
created hot spots within the film under an electrical bias and accelerated breakdown,
reducing energy density.
24,29
Song, Tang and coworkers additionally showed that
modified high aspect ratio BaTiO
3
nanofibers can also enhance the dielectric constant,
breakdown strength, and energy density in high permittivity PVDF-based composite
systems.
49,50

As an extension of their earlier work (vide supra) Tang and coworkers employed a
double-layer core/shell-like structure to improve the interfacial interaction in
PEN/BaTiO
3
nanocomposites.
51
BaTiO
3
nanocrystals were amine terminated through a
silane coupling agent then grafted by electrophilic addition to a nitrile-based
hyperbranched bifunctional copper phthalocyanine forming a shell approximately 25-30
nm thick. The bifunctional nitrile groups allowed for strong covalent bonding with the
amine functionalized BaTiO
3
, and subsequent interaction with PEN, confirmed by FT-IR.
The polymeric interaction with the phthalocyanine-modified BaTiO
3
surface displayed a
continual decrease in the infrared CN stretching frequency at 2230 cm
-1
as the polymer
concentration increased.
51
Despite a favorable interaction, the dielectric constant and loss
tangent decreased 16% over unmodified BaTiO
3
, and additionally showed significant
temperature dependent dielectric phenomena. The lower dielectric constant was a result
14

of the significant surface modification, which reduced the PEN-BaTiO
3
interaction within
the interfacial zone, consistent with other reports on modified BaTiO
3
.
42,43,52

Using isopropyl tris(N-amino-ethyl aminoethyl)titanate, a novel inorganic coupling agent,
to modify BaTiO
3
, Wang and coworkers prepared composites with polyethersulfone
(PES) under high temperature conditions.
53
The composites were stable to above 500 °C
at all loadings, and additionally displayed very little dielectric temperature dependence up
to 150 °C. At 50 vol% BaTiO
3
, the permittivity reached ɛ
r
= 24.6, increasing more than
five times that of pure PES. The permittivity was not compared against effective medium
models, but is consistent with reports from Calame and coworkers showing near
exponential permittivity increases when filler concentrations rise above 30 vol%.
54,55

Direct mixing techniques have shown to be effective in studying the interaction of
nanocrystalline BaTiO
3
on polymeric systems. The hydroxylated surface of BaTiO
3
can
contribute to significant hydrogen bonding within the interfacial zone, altering the
dielectric response of the composite in an unpredictable fashion. Significant surface
modification, on the other hand, can potentially mitigate the interfacial effect and
improve nanocrystal dispersion without adversely modifying the inherent polymer
properties, providing the opportunity to simplify and better control an already complex
and unpredictable system. At this point, there is no universal correlation to the composite
dielectric performance and the range of interactions studied.

15

1.3 Nanocomposite Preparation via In Situ Polymerization
To gain an added level of control over direct mixing techniques which use pre-formed
polymers dissolved in solution, researchers have studied the effects of in situ
polymerization in the presence of BaTiO
3
nanocrystals.
56
For in situ polymerization, a
nanocrystalline suspension is directly added to a separate solution containing
monomer(s). The mixture is homogenized to disperse the nanocrystals and then treated
under high temperatures or with reactive species (e.g., a catalyst or initiator) to initiate
polymeric chain growth.
56
By homogenizing the nanocrystal solution with the
monomer(s) prior to polymerization, it is possible to enhance the interaction and
potentially use the nanocrystal as a vehicle to chemically graft to, or to become part of,
the polymer matrix. In situ polymerization additionally provides access to traditionally
insoluble or non-processable polymers, like polyimides and epoxies, that have excellent
dielectric and physical characteristics.
28,57

1.3.1 In Situ Polymerization in the Presence of Unmodified BaTiO
3

Epoxy resins are among the most widely studied materials for in situ polymerization in
the presence of BaTiO
3
. Epoxy resins are a class of highly reactive oligomers containing
epoxide functional groups. Alone, epoxy resins have poor mechanical, thermal, and
dielectric properties; however, the properties can be drastically improved and controlled
by chemically cross-linking the resins with a wide range of co-reactants including
amines, acids, alcohols, or thiols.
57
Epoxy-based polymers have been used in a wide
range of applications, and are currently the primary matrix material for BaTiO
3
based
16

embedded capacitors.
41
In this role, a small amount of epoxy acts as a binder for micron-
sized BaTiO
3
resulting in rigid and non-flexible composites. Patsidis and coworkers have
sought to overcome these limitations and have successfully prepared epoxy/BaTiO
3

composites at different loadings using commercial BaTiO
3
of 50 nm and 2 μm in
diameter. To prepare composites, BaTiO
3
particles of varying concentrations were
dispersed in isopropyl alcohol followed by addition of a commercial epoxy resin,
Araldite LY 1564. The mixture was stirred for 1 h and the curing agent, Aradur HY2954,
was added to the solution and cast into a mold and cured.
58
Through the use of XRD,
they confirmed that the nanocrystalline BaTiO
3
was a mixture of the tetragonal and cubic
phases whereas the micrometer sized particles were highly tetragonal. Both crystallites
underwent a complete phase change above 130 °C to the paraelectric cubic phase.
Frequency dependant dielectric measurements were performed at 30 and 150 °C for both
sets of particles at 5, 7, and 10 wt% particle loading. Under all concentrations and
temperatures studied, it was clear that the larger BaTiO
3
crystallites displayed higher

Scheme 1.3. Proposed interaction of PAA oligomers with the surface of BaTiO
3
. Modified from
Reference 23.
17

permittivities. However, as the concentration of filler increased, the permittivity of the
micro- and nano-sized BaTiO
3
did not follow effective medium models, which predict a
systematic increase in permittivity as the concentration of particles increases.
38
Instead,
the relative permittivity reached a maximum at 7 wt% BaTiO
3
and fell sharply at 10 wt%.
This reduction has been reported in epoxy/oxide systems and was determined to be a
result of unaligned polar domains, voids, and the presence of filler agglomerates at higher
concentrations.
41,58,59

Polyimides are another class of polymers which have gained significant attention as a
matrix material for BaTiO
3
. Unlike epoxies which require an initiator and often complex
proprietary resin systems, polyimides are highly versatile classes of condensation
polymers. Polyimides can be prepared by the spontaneous reaction of any diamine and
dianhydride species in polar aprotic solvents (e.g., DMF, DMSO, or N,N’-
dimethylacetamide (DMAc)) forming a solution processable reactive oligomer, called a
polyamic acid (PAA). The formation of PAA and its propsed interaction with the surface
of a BaTiO
3
nanocrystal is depicted in Scheme 1.3. The PAA can be cast into films or
spun onto substrates, after which time the film can be heated, removing solvent and
driving a condensation reaction that irreversibly sets the polymer.
60,61
The imidized
polymer, depending on the backbone rigidity, can be a solution processable thermoplastic
with low melting points or an unprocessable thermoset with high thermal stability.
62,63

Polyimides are being explored as an alternative polymer component for nanocomposites
because of their versatility, high thermal and chemical stability, low dielectric loss, and
high breakdown strength.
56,64–66
Dang and coworkers prepared a polyimide/BaTiO
3

18

nanocomposite system by in-situ polymerization. Unmodified BaTiO
3
nanocrystals
approximately 100 nm in diameter were prepared by a hydrothermal technique, dispersed
in DMAc, and added to a solution containing 4,4’-oxydianiline (ODA), followed by slow
addition of pyromellitic dianhydride (PMDA). Low molecular weight PAA oligomers
formed during the polymerization were postulated to aid in dispersion through a
hydrogen bonding interaction at the interfacial zone (Scheme 1.3), forming a core/shell
like structure.
23
The composites were stable to 500 °C and the relative permittivity
increased linearly up to 40 vol% loading, reaching ɛ
r
= 18. Despite the increase in
permittivity, the dielectric breakdown strength decreased by 40% with only 5 vol%
BaTiO
3
, indicating the onset of percolation. The steep decrease in breakdown voltage
allowed for only marginal increase in the calculated energy density at 40 vol% BaTiO
3
.
Feng and coworkers also prepared nanocomposites with 100 nm unmodified BaTiO
3
in
the same polymer system and found similar trends in decreasing breakdown strength at
low filler loadings.
67

1.3.2 In Situ Polymerization in the Presence of Modified BaTiO
3

In a similar approach to the direct addition systems (vide supra), modification of the
nanocrystal surface can be used to further enhance the interaction of BaTiO
3
to
potentially aid in dispersion and improve dielectric performance. Choudhury and
coworkers utilized hydrogen peroxide treatment of BaTiO
3
to successfully prepare
polyetherimide/BaTiO
3
(PEI/BaTiO
3
) nanocomposites.
68
1,3-phenylenediamine and
4,4’-(4,4’-isopropylidenediphenoxy) bis(phthalic anhydride) were mixed in N-methyl-2-
19

pyrrolidone (NMP), forming a viscous PAA solution in the presence of hydroxylated 70
nm BaTiO
3
nanocrystals. The solutions were spun cast and thermally imidized to 250 °C,
forming films approximately 150 nm thick. The strong interaction between the hydroxyl-
functionalized BaTiO
3
and the PEI matrix greatly enhanced the particle dispersion and
displayed no effect on the polymer, which completely imidized. X-ray diffraction
analysis indicated that the nanocrystals remained phase pure and were largely unchanged
by the processing conditions, whereas the polymer crystallinity decreased in the presence
of BaTiO
3
. Dielectric analysis of PEI/BaTiO
3
showed a marked increase in the relative
permittivity with a maximum ε
r
= 40 at 50 vol% BaTiO
3
. The higher than expected
permittivity was postulated to be a result of localized field interactions of neighboring
BaTiO
3
nanocrystals within the matrix, due to the enhanced interaction from
hydroxylation.
41,68

Through another peroxide mediated approach, Xie and coworkers added amine
functionality to 100 nm BaTiO
3
with a silane coupling agent, 3-aminopropyl-

Scheme 1.4. Schematic diagram illustrating the process of ATRP from the surface of BaTiO
3
.
Modified from reference 71.
20

triethoxysilane, to prepare polyimide-based nanocomposites using the PMDA/ODA
system (vide supra).
69
The presence of amine groups at the surface of the nanocrystals
aided in nanoparticle dispersion and additionally provided a direct cross-linking site that
the PAA oligomer could interact with. The enhanced interaction increased the thermal
stability beyond that of pure polymer and eliminated the effects of catalytic thermal
oxidation, commonly present in polyimide-oxide composites.
65
XRD analysis of the films
revealed that the presence of BaTiO
3
reduced the polyimide packing density as evidenced
by a peak shift in the amorphous region. Due to the enhanced interaction at the interface
as a result of surface modification, the composite displayed stable relative permittivities
above ε
r
= 34 at 50 vol% BaTiO
3
across a range of frequencies and temperatures. The
dielectric permittivity in this system was improved by 88% over the composites prepared
by Dang and coworkers using unmodified BaTiO
3
of the same size.
23
While both
systems provided an interaction within the interfacial zone, the amine moiety from Xie
and coworkers created a covalent nucleating site by which the polymer could grow off of,
explaining the dielectric response.
Using a two-step approach involving hydrogen peroxide activation, Siddabattuni and
coworkers studied the effect of controlled interfacial chemistry on the properties of
epoxy/BaTiO
3
nanocomposites.
70
By pre-treating BaTiO
3
with hydrogen peroxide they
were then able to effectively graft a dense layer of 2-aminoethyl dihydrogen phosphate
(AEP) to the surface. The high concentrations of amino groups on the surface were then
capable of more efficiently reacting, and cross-linking with, the epoxy polymer matrix.
The resulting interfacial interaction reduced the polymer free volume and raised the
21

composite thermal stability, improving electrical breakdown resistance. The as-prepared
composite had a higher relative permittivity (ε
r
= 6.3 at 5 vol%), a lower dielectric loss
(tan δ = 0.017), and a higher breakdown strength (406 V μm
–1
) than any of their controls.
The improved breakdown resistance was thought to be a result of reduced free charge
carriers, but it was not clear as to whether the concentration of free carriers was
physically reduced or if the carriers had a lower mobility within the composite.
70

The group of Jiang and coworkers developed a complex process involving hydrogen
peroxide treatment and direct surface modification of BaTiO
3
to prepare PMMA
nanocomposites via in situ atom transfer radical polymerization, depicted in Scheme 1.4
(ATRP).
71
By functionalizing the surface of BaTiO
3
with 3-aminopropyl triethoxysilane,
they were able to take advantage of the amine-rich surface and graft α-bromoisobutyryl
bromide using it as an initiating site for ATRP. A controlled amount of initiator and
methyl methacrylate (MMA) were added to the brominated BaTiO
3
solution and heated,

Figure 1.2. TEM images of PMMA/BaTiO
3
displaying shell thicknesses of (a) 7.5 nm, (b) 10
nm, (c) 15 nm, and (d) 17 nm. Modified from reference 71.
22

forming a controlled core/shell-like structure around each BaTiO
3
nanocrystal with
thicknesses ranging from 7.5-17 nm, shown in Figure1.2 (thickness dependent on
MMA:BaTiO
3
ratio). As the thickness of the shell increased, the composite became more
dispersable but also displayed a lower relative permittivity, reaching optimum conditions
with a 10 nm shell. At this composition the permittivity reached ε
r
= 20 and losses were
below tan δ = 0.06 at all frequencies and temperatures. Thinner core/shell structures
were not stable in solution and could not be studied. In more recent work; however, they
were able to graft S-1-ethyl-S'-(α,α'-dimethyl- α''-acetic acid)trithiocarbonate (EDMAT),
a known reagent to facilitate reversible addition-fragmentation chain-transfer (RAFT)
polymerization of polystyrene, onto the surface of amine modified BaTiO
3
nanocrystals
to prepare polystyrene/BaTiO
3
nanocomposites. Their composites, prepared similarly to
the PMMA/BaTiO
3
composites, displayed a controllable core-shell structure ranging in
thickness from 7-12 nm. The use of polystyrene proved beneficial, by increasing the
relative permittivity to ε
r
= 24.51 with a 7 nm shell, a more than 20% increase over their
previously published work. Additionally, the reported dielectric loss reduced more than
50% falling below tan δ = 0.03 at all frequencies and concentrations.
72
The lower
permittivity and higher loss reported in the PMMA/BaTiO
3
composite systems was likely
a result of the presence of charged bromine species and polarizable MMA moieties which
prevented the processing shells thinner than 10 nm. Polystyrene on the other hand is
weakly polarizable and its grafting agent was non-ionic, allowing for lower losses and
more stable dispersions with a thinner shell thickness.
23

Through a novel UV curing sol-gel based grafting technique, Chon and coworkers were
able to successfully prepare nanocomposites with 80 nm and 4 μm BaTiO
3
crystallites at
a range of loadings.
73
The composite system was based on highly polarizable silane and
bisphenol A moieties which aided in nanocrystal dispersion and improved dielectric
performance over unmodified materials. The surface modification and composite
formation process is highlighted in Scheme 1.5, with the final curing step taking place

Scheme 1.5. Schematic representation of BaTiO
3
surface modification, grafting, and curing in
the formation of bisphenol A based nanocomposite. Modified from reference 73.
24

after spin casting the device onto tin doped indium oxide (ITO) coated glass substrates.
Compositions containing less than 14 vol% BaTiO
3
could be peeled off the substrate,
showing high mechanical strength and flexibility. At 50 vol% loading of modified
BaTiO
3
, the relative permittivity was ε
r
= 25 and 62 for the 80 nm and 4 μm particles,
respectively, representing a 9 and 13 % increase over unmodified BaTiO
3
at the same
loading. The maximized dielectric strength was reported to be 263 V μm
–1
for the 80 nm
BaTiO
3
nanocrystals, but neither the concentration of this optimized device nor the data
for the 4 μm crystals was reported.
73

In situ polymerization provides an added level of control over direct addition-based
composite systems, and allows the use of polymers which would otherwise be
inaccessible. Additionally, adding reactive functional groups at the surface of BaTiO
3

which participate in polymerization can improve dielectric permittivity, loss, breakdown
strength, and stability. While dispersion and interfacial effects are attributed to the
improved properties, there is still a lot which remains unknown. For in situ systems it is
difficult to determine the properties of polymer (i.e., molecular weight or exact structure),
and whether it is in the same form as if it was prepared without the presence of BaTiO
3
.
When considering the high surface-to-volume ratio and reactivity of BaTiO
3

nanocrystals, it is reasonable to surmise that chemical reactions at the surface can
severely alter the polymerization process and could even change the properties of the
nanocrystal, affecting the perceived dielectric response.
74

25

1.4 Conclusions and Perspective
The use of nanocrystalline BaTiO
3
in polymer nanocomposite dielectrics provides the
opportunity to significantly enhance the dielectric response of a material. Recent
advancements have displayed the necessity of improved dispersion in harnessing
dielectrics that display solution processability, flexibility, high permittivity, and superior
dielectric strength. Although it is clear that the dispersion of nanocrystalline BaTiO
3
is
dependent upon interfacial interactions, it still remains unclear as to how these effects can
be systematically controlled. Surface modification allows us to qualitatively predict
and/or enhance the interaction of a nanocrystal with its environment in order to form
homogeneous solutions, but little attention is given to how these modifications alter the
electronic properties of a nanocrystal before it is introduced into a polymer matrix. In
situ systems add another level of complexity because the individual nature of BaTiO
3
and
the polymer are largely unknown. Understanding the chemically induced effects of the
interfacial zone on nanocrystalline BaTiO
3
-based polymer nanocomposites will provide a
means by which we can begin to determine why BaTiO
3
can cause atypical concentration
and size dependence, display higher and lower permittivities, and induce crystallinity or
disorder. It is clear from the studies presented that these interactions play a pivotal role
in determining device performance, and that the individual properties of the nanocrystal
and polymer must not be the only consideration when designing high energy density
capacitors. By taking a more systematic approach to composite formation we can begin
to make substantial, less incremental improvements.
26

1.5 References
(1) Goldschmidt, V. M. Skrifter Norske Videnskaps-Akad. Oslo. I. Matemal.-Naturv.
Klasse 1927, 8, 7–156.

(2) Newnham, R. E.; Cross, L. E. MRS Bull. 2011, 30, 845–848.

(3) Lines, M. E. Principles and applications of ferroelectrics and related materials;
Oxford classic texts in the physical sciences; Clarendon Press ; Oxford University
Press: Oxford : New York, 2001.

(4) Merz, W. Phys. Rev. 1949, 75, 687–687.

(5) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Small 2008, 4, 2102–2106.

(6) Brutchey, R. L.; Morse, D. E. Angew. Chem. Int. Ed. 2006, 45, 6564–6566.

(7) Brutchey, R. L.; Cheng, G.; Gu, Q.; Morse, D. E. Adv. Mater. 2008, 20, 1029–
1033.

(8) Mao, Y.; Banerjee, S.; Wong, S. S. J. Am. Chem. Soc. 2003, 125, 15718–15719.

(9) Mao, Y.; Banerjee, S.; Wong, S. S. Chem. Commun. 2003, 408–409.

(10) O’Brien, S.; Brus, L.; Murray, C. B. J. Am. Chem. Soc. 2001, 123, 12085–12086.

(11) Gaskins, B. C.; Lannutti, J. J. J. Mater. Res. 2011, 11, 1953–1959.

(12) Niederberger, M.; Pinna, N.; Polleux, J.; Antonietti, M. Angew. Chem. Int. Ed.
2004, 43, 2270–2273.

(13) Spaldin, N. A. Science 2004, 304, 1606–1607.

(14) Buscaglia, V.; Buscaglia, M. T.; Viviani, M.; Mitoseriu, L.; Nanni, P.; Trefiletti,
V.; Piaggio, P.; Gregora, I.; Ostapchuk, T.; Pokorný, J.; Petzelt, J. J. Eur. Ceram.
Soc. 2006, 26, 2889–2898.

(15) Smith, M. B.; Page, K.; Siegrist, T.; Redmond, P. L.; Walter, E. C.; Seshadri, R.;
Brus, L. E.; Steigerwald, M. L. J. Am. Chem. Soc. 2008, 130, 6955–6963.

(16) Zhao, Z.; Buscaglia, V.; Viviani, M.; Buscaglia, M.; Mitoseriu, L.; Testino, A.;
Nygren, M.; Johnsson, M.; Nanni, P. Phys. Rev. B. 2004, 70.

27

(17) Tsurumi, T.; Sekine, T.; Kakemoto, H.; Hoshina, T.; Nam, S.-M.; Yasuno, H.;
Wada, S. J. Am. Ceram. Soc. 2006, 89, 1337–1341.

(18) Zembilgotov, A. G.; Pertsev, N. A.; Waser, R. J. Appl. Phys. 2005, 97, 114315.

(19) Takeuchi, T.; Tabuchi, M.; Ado, K.; Honjo, K.; Nakamura, O.; Kageyama, H.;
Suyama, Y.; Ohtori, N.; Nagasawa, M. J. Mater. Sci. 1997, 32, 4053–4060.

(20) Bansal, V.; Poddar, P.; Ahmad, A.; Sastry, M. J. Am. Chem. Soc. 2006, 128,
11958–11963.

(21) Yen, F.-S.; Chang, C. T.; Chang, Y.-H. J. Am. Ceram. Soc. 1990, 73, 3422–3427.

(22) Kim, P.; Doss, N. M.; Tillotson, J. P.; Hotchkiss, P. J.; Pan, M.-J.; Marder, S. R.;
Li, J.; Calame, J. P.; Perry, J. W. ACS Nano 2009, 3, 2581–2592.

(23) Dang, Z.-M.; Lin, Y.-Q.; Xu, H.-P.; Shi, C.-Y.; Li, S.-T.; Bai, J. Adv. Funct.
Mater. 2008, 18, 1509–1517.

(24) Barber, P.; Balasubramanian, S.; Anguchamy, Y.; Gong, S.; Wibowo, A.; Gao,
H.; Ploehn, H. J.; Zur Loye, H.-C. Materials 2009, 2, 1697–1733.

(25) Ash, B. J.; Siegel, R. W.; Schadler, L. S. J. Polym. Sci. Pol. Phys. 2004, 42,
4371–4383.

(26) Khalil, M. S. IEEE T. Dielect. El. In. 2000, 7, 261–268.

(27) Ma, D.; Hugener, T. A.; Siegel, R. W.; Christerson, A.; Mårtensson, E.; Önneby,
C.; Schadler, L. S. Nanotechnology 2005, 16, 724–731.

(28) Nanocomposites: preparation, properties, and performance; Nova Science
Publishers: New York, 2008.

(29) Calebrese, C.; Hui, L.; Schadler, L.; Nelson, J. IEEE T. Dielect. El. In. 2011, 18,
938–945.

(30) Tanaka, T.; Montanari, G. C.; Mulhaupt, R. IEEE T. Dielect. El. In. 2004, 11,
763–784.

(31) Drobny, J. G.; Rapra Technology Limited Fluoroplastics; Rapra Technology Ltd.:
Shawbury, Shrewsbury, Shropshire, U.K., 2006.

(32) Wang, Y.; Zhou, X.; Chen, Q.; Chu, B.; Zhang, Q. IEEE T. Dielect. El. In. 2010,
17, 1036–1042.
28

(33) Mendes, S. F.; Costa, C. M.; Caparros, C.; Sencadas, V.; Lanceros-Méndez, S. J.
Mater. Sci. 2011, 47, 1378–1388.

(34) Dang, Z.-M.; Wang, H.-Y.; Peng, B.; Nan, C.-W. J. Electroceram. 2007, 21, 381–
384.

(35) Mao, Y. P.; Mao, S. Y.; Ye, Z.-G.; Xie, Z. X.; Zheng, L. S. J. Appl. Phys. 2010,
108, 014102.

(36) Kobayashi, Y.; Tanase, T.; Tabata, T.; Miwa, T.; Konno, M. J. Eur. Ceram. Soc.
2008, 28, 117–122.

(37) Pant, H.; Patra, M.; Verma, A.; Vadera, S.; Kumar, N. Acta Mater. 2006, 54,
3163–3169.

(38) Tinga, W. R. J. Appl. Phys. 1973, 44, 3897.

(39) Kim, D.-W.; Lee, D.-H.; Kim, B.-K.; Je, H.-J.; Park, J.-G. Macromol. Rapid
Comm. 2006, 27, 1821–1825.

(40) Tang, H.; Zhong, J.; Yang, J.; Ma, Z.; Liu, X. J. Electron. Mater. 2010, 40, 141–
148.

(41) Cho, S.-D.; Lee, S.-Y.; Hyun, J.-G.; Paik, K.-W. J. Mater. Sci.-Mater. El. 2005,
16, 77–84.

(42) Zhou, T.; Zha, J.-W.; Cui, R.-Y.; Fan, B.-H.; Yuan, J.-K.; Dang, Z.-M. ACS. Appl.
Mater. Interfaces 2011, 3, 2184–2188.

(43) Almadhoun, M. N.; Bhansali, U. S.; Alshareef, H. N. J. Mater. Chem. 2012, 22,
11196.

(44) Jung, H. M.; Kang, J.-H.; Yang, S. Y.; Won, J. C.; Kim, Y. S. Chem. Mater. 2010,
22, 450–456.

(45) Chu, L. W.; Prakash, K. N.; Tsai, M.-T.; Lin, I.-N. J. Eur. Ceram. Soc. 2008, 28,
1205–1212.

(46) Neouze, M.-A.; Schubert, U. Montash. Chem. 2008, 139, 183–195.

(47) Kim, P.; Jones, S. C.; Hotchkiss, P. J.; Haddock, J. N.; Kippelen, B.; Marder, S.
R.; Perry, J. W. Adv. Mater. 2007, 19, 1001–1005.

29

(48) Li, J.; Claude, J.; Norena-Franco, L. E.; Seok, S. I.; Wang, Q. Chem. Mater. 2008,
20, 6304–6306.

(49) Tang, H.; Lin, Y.; Sodano, H. A. Goulbourne, N. C.; Ounaies, Z., Eds.; 2012; pp.
834206–834206–8.

(50) Song, Y.; Shen, Y.; Liu, H.; Lin, Y.; Li, M.; Nan, C.-W. J. Mater. Chem. 2012,
22, 16491.

(51) Tang, H.; Ma, Z.; Zhong, J.; Yang, J.; Zhao, R.; Liu, X. Colloid Surface A. 2011,
384, 311–317.

(52) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Langmuir 2010, 26, 5067–5071.

(53) Wang, F. J.; Li, W.; Xue, M. S.; Yao, J. P.; Lu, J. S. Compos. Part B.-Eng. 2011,
42, 87–91.

(54) An, A. L.; Boggs, S. A.; Calame, J. IEEE Int. Symp. El. In. 2006, 466–469.

(55) An, A. L.; Boggs, S. A.; Calame, J. IEEE Electr. Insul. M. 2008, 24, 5–10.

(56) Chen, B.; Evans, J. R. G.; Greenwell, H. C.; Boulet, P.; Coveney, P. V.; Bowden,
A. A.; Whiting, A. Chem. Soc. Rev. 2008, 37, 568.

(57) Friedrich, K. Polymer composites: from nano-to-macro-scale; Springer: New
York, 2005.

(58) Patsidis, A. C.; Kalaitzidou, K.; Psarras, G. C. Mater. Chem. Phys. 2012, 135,
798–805.

(59) Tomer, V.; Polizos, G.; Manias, E.; Randall, C. A. J. Appl. Phys. 2010, 108,
074116.

(60) Sroog, C. E.; Endrey, A. L.; Abramo, S. V.; Berr, C. E.; Edwards, W. M.; Olivier,
K. L. J. Polym. Sci. Part A. 1965, 3, 1373–1390.

(61) Sroog, C. E. J. Polym. Sci. Macromol. Rev. 1976, 11, 161–208.

(62) Mittal, K. L. Polyimides and other high temperature polymers: synthesis,
characterization and applications; VSP: Savannah, GA, 2007.

(63) Tamai, S.; Yamaguchi, A.; Ohta, M. Polymer 1996, 37, 3683–3692.

30

(64) Fragiadakis, D.; Logakis, E.; Pissis, P.; Kramarenko, V. Y.; Shantalii, T. A.;
Karpova, I. L.; Dragan, K. S.; Privalko, E. G.; Usenko, A. A.; Privalko, V. P. J.
Phys. Conf. Ser. 2005, 10, 139–142.

(65) Lin, B.; Liu, H.; Zhang, S.; Yuan, C. J. Solid State Chem. 2004, 177, 3849–3852.

(66) Murugaraj, P.; Mainwaring, D.; Mora-Huertas, N. J. Appl. Phys. 2005, 98,
054304.

(67) Yu Feng; Jinghua Yin; Minghua Chen; Xiaoxu Liu; Guang Li IEEE, 2011; pp.
226–229.

(68) Choudhury, A. Mater. Chem. Phys. 2010, 121, 280–285.

(69) Xie, S.-H.; Zhu, B.-K.; Wei, X.-Z.; Xu, Z.-K.; Xu, Y.-Y. Compos. Part A.-Appl.
S. 2005, 36, 1152–1157.

(70) Siddabattuni, S.; Schuman, T. P.; Dogan, F. Mater. Sci. Eng. B. 2011, 176, 1422–
1429.

(71) Xie, L.; Huang, X.; Wu, C.; Jiang, P. J. Mater. Chem. 2011, 21, 5897.

(72) Yang, K.; Huang, X.; Xie, L.; Wu, C.; Jiang, P.; Tanaka, T. Macromol. Rapid
Comm. 2012, 33, 1921–1926.

(73) Chon, J.; Ye, S.; Cha, K. J.; Lee, S. C.; Koo, Y. S.; Jung, J. H.; Kwon, Y. K.
Chem. Mater. 2010, 22, 5445–5452.

(74) Fernández-García, M.; Martínez-Arias, A.; Hanson, J. C.; Rodriguez, J. A. Chem.
Rev. 2004, 104, 4063–4104.

31

Chapter 2. Low-Temperature Synthesis of Solid-Solution Ba
x
Sr
1-x
TiO
3

Nanocrystals*

*Published in J. Mater. Chem. 2010, 20, 5074.
2.1 Introduction
Perovskites, such as alkaline earth metal titanates (ATiO
3
, where A = Ca, Sr, Ba), can
exhibit a wide range of technologically important dielectric, ferroelectric,
piezoelectric, and pyroelectric properties, in addition to having a positive temperature
coefficient of resistivity.
1,2
These perovskite structures can be doped with small
amounts of iso- and aliovalent cations, while the addition of stoichiometric amounts
of isovalent cations will form true solid solutions.
3
It is known that altering the
composition via doping or solid solution formation has a profound effect on the
properties of the perovskite material.
4,5
Moreover, there has been a great deal of
interest in studying the properties of perovskites at the nanoscale;
6-9
however, there
have not been many studies on the effects of both size and composition on the
properties of well-defined (and unsintered) perovskite nanocrystals. This may be a
direct result of the inherent difficulty in synthesizing perovskite nanocrystals that are
well defined, compositionally complex, and phase pure; that is, an accurate control
over stoichiometry, precursor reactivity, and conditions necessary for crystallization
(i.e., temperature, pressure, and pH) are all critically important. Even with careful
synthetic consideration, phase segregation into two or more crystalline or amorphous
phases (e.g., TiO
2
, BaCO
3
, and/or Ba
2
TiO
4
for BaTiO
3
) can occur as a result of these
complicating factors.
10,11

32

Titanates are traditionally prepared via high temperature (>1100 ˚C) solid-state
reactions between TiO
2
and alkaline earth metal carbonates (e.g., BaCO
3
or SrCO
3
),
which yield large, micron-sized crystal grains with wide ranges in size and shape.
3

Deposition techniques such as metal-organic chemical vapor deposition, pulsed laser
ablation, sputtering, and molecular beam epitaxy provide a high level of structural
control; however, these methods are energy intensive and expensive.
12
Thus,
inexpensive solution-phase routes (e.g., sol-gel, hydrothermal, solvothermal, or
micellar) are being utilized in order to gain control over perovskite size and
morphology. In the past 10 years, significant strides have been made in the solution-
phase syntheses of well-defined ternary perovskite nanocrystals, such as BaTiO
3
and
SrTiO
3
.
13-16
Unfortunately, the same level of nanostructural control over the synthesis
of more compositionally complex solid solution perovskites is still lacking. For
example, Ba
x
Sr
1-x
TiO
3
is an important solid-solution perovskite that possesses unique
dielectric properties that are dependent upon composition. The vast majority of
solution-phase routes to nanostructured Ba
x
Sr
1-x
TiO
3
result in relatively large particles
≥50 nm in size.
17-23
Notably, several reports of well-defined, sub-50 nm Ba
x
Sr
1-x
TiO
3

nanocrystals have recently appeared in the literature. Niederberger et al. first reported
the synthesis of small, 4-nm Ba
0.5
Sr
0.5
TiO
3
nanocrystals using a nonaqueous
method.
24
Su et al. subsequently reported the synthesis of 10-50 nm Ba
x
Sr
1-x
TiO
3

nanocrystals of variable composition using a reverse micelle route.
25
Likewise, Wei
et al. described the synthesis of 15-36 nm Ba
x
Sr
1-x
TiO
3
nanocrystals of variable
composition using a solvothermal technique.
26
Each of these methods requires either
33

strongly alkaline conditions,
25
reaction temperatures of 200 ˚C,
24
or combinations
thereof for particle formation and crystallization.
26
Thus, the challenge still exists to
make small, well-defined Ba
x
Sr
1-x
TiO
3
solid-solution nanocrystals using scalable and
ultrabenign reaction conditions (i.e., low temperature, neutral pH, ambient
pressure).
27
The vapor diffusion sol-gel route was recently reported as a method of synthesizing 6-nm
BaTiO
3
nanocrystals at room temperature, ambient pressure, and near-neutral pH.
28
The
success of this method stems from the fact that a bimetallic alkoxide is hydrolyzed in a
kinetically controlled manner via the introduction of water vapor. As a result, particle
nucleation and growth occur at room temperature and a post-synthesis annealing step for
crystallization via solid-solid diffusion is not required. Following this initial report, it
was subsequently discovered that small concentrations (0.05–0.17 at%) of aliovalent La
3+

dopants can be controllably introduced into the 6-nm BaTiO
3
nanocrystals during
synthesis under the same benign conditions.
29
Herein, we demonstrate that by tuning the
reaction temperature, the vapor diffusion sol-gel method can be extended to make yet
more compositionally complex Ba
x
Sr
1-x
TiO
3
solid-solution nanocrystals that are small
(sub-15 nm), well defined, and synthesized at low temperatures, ambient pressure and
neutral pH. The synthesis of small, well-defined Ba
x
Sr
1-x
TiO
3
nanocrystals allowed their
dielectric properties to be studied as a function of composition at the nanoscale. These
results represent the lowest temperature at which Ba
x
Sr
1-x
TiO
3
nanocrystals have been
synthesized, and the highest dielectric constant reported thus far for a sub-30 nm Ba
x
Sr
1-
x
TiO
3
nanocrystal.
34

2.2 Experimental Section
2.2.1 General Procedures
All manipulations were performed under a nitrogen atmosphere using standard Schlenk
techniques. Air-free solvents were used throughout. Polyvinyl alcohol (98-99%
hydrolyzed; Alfa Aesar, Inc.), barium titanium bimetallic alkoxide
(BaTi[OCH
2
CH(CH
3
)OCH
3
]
6
; 0.5 M in n-butanol/3-methoxypropanol; Gelest, Inc.), and
strontium titanium bimetallic alkoxide (SrTi[OCH
2
CH(CH
3
)OCH
3
]
6
; 0.7 M in n-
butanol/3-methoxypropanol; Gelest, Inc.) were used as received.
2.2.2 Synthesis of Ba
x
Sr
1-x
TiO
3
(0 ≤ x ≤ 1) Nanocrystals
Deionized water (20 mL) and 10 mL of a separate 0.43 M solution of barium titanium
bimetallic alkoxide and strontium titanium bimetallic alkoxide (molar ratio dependent
upon desired composition) was placed into an enclosed chamber under a nitrogen
atmosphere. The solution containing the bimetallic alkoxides was exposed to water
vapor for 24 h at 20
º
C for compositions where x > 0.5. After 24 h, the water source was
removed and the particles were allowed to age under nitrogen at 37 ˚C for 24 h. The
resulting off-white gel was collected and rinsed in ethanol (3  10 mL) and dried in vacuo
(25 ˚C, 0.05 mm Hg) for 24 h to yield a nanoparticulate powder of Ba
x
Sr
1-x
TiO
3
(1.0 g).
For compositions where x ≤ 0.5, the reaction was performed at 80
º
C for 48 hours with no
aging required and otherwise identical conditions.

35

2.2.3 Material Characterization
Transmission electron microscopy (TEM) analysis was performed on a JEOL JEM-2100
microscope at an operating voltage of 120 kV, which was equipped with a Gatan Orius
CCD camera. Samples for TEM analysis were prepared by drop casting a suspension of
the nanocrystals in methanol onto ultrathin carbon film supported on 400 mesh copper
grids (Ted Pella, Inc.). BET measurements were performed on a Nova 2200e surface
area and pore size analyzer (Quantachrome Instruments, Inc.). Powder X-ray diffraction
(XRD) was performed on a Rigaku Ultima IV diffractometer using a Cu Kα (λ = 1.54 Å)
radiation. Silicon powder was used as the internal standard. For XRD samples,
suspensions of Ba
x
Sr
1-x
TiO
3
(0 ≤ x ≤ 1) nanocrystals in methanol were deposited on glass
substrates and dried at room temperature. Elemental analysis was performed via
inductively coupled plasma-optical emission spectrometry (ICP-OES) at the
Microanalysis Laboratory at the University of Illinois at Urbana-Champaign.
2.2.4 Dielectric Characterization
For ensemble dielectric studies, a 13-mm diameter pellet was prepared by grinding 200
mg of the Ba
x
Sr
1-x
TiO
3
nanocrystalline powder with 1 mL of an aqueous solution of
polyvinyl alcohol (PVA; 1 mg/mL). The slurry was allowed to dry overnight in a
nitrogen atmosphere. The dried powder was pressed with an applied load of 6 metric
tons in vacuo (10 mm Hg). After pressing, the pellet was annealed under nitrogen at 150

ᵒ
C for 3 h. Colloidal silver paint (Ted Pella, Inc.) was applied to both sides before
annealing the pellet at 100
ᵒ
C for 1 h under nitrogen. Capacitance and loss tangents were
36

measured on an Agilent 4294A Impedance Analyzer with a frequency sweep of 1 kHz to
2 MHz. Repeated measurements were taken on two independently prepared samples of
each composition over the course of several weeks, and the data were shown to be
consistent between the different sample preparations and do not change with time.
2.3 Results and Discussion
2.3.1 Nanocrystal Synthesis and Characterization
Solid solutions of Ba
x
Sr
1-x
TiO
3
nanocrystals (x = 0, 0.14, 0.27, 0.33, 0.43, 0.48, 0.63,
0.69, 0.77, and 1.00) were prepared by a modified vapor diffusion sol-gel technique,
whereby a bimetallic BaTi[OCH
2
CH(CH
3
)OCH
3
]
6
alkoxide solution (0.5 M in n-
BuOH/2-methoxypropanol) was mixed with a bimetallic SrTi[OCH
2
CH(CH
3
)OCH
3
]
6

alkoxide solution (0.7 M in n-BuOH/2-methoxypropanol) and allowed to slowly react
with water vapor over 24 h. By varying the ratio of the two bimetallic alkoxides,
solid-solution Ba
x
Sr
1-x
TiO
3
nanocrystals over the entire composition range were

Figure 2.1. (a) Powder X-ray diffraction pattern of the cubic SrTiO
3
nanocrystals. (b) Low-
resolution TEM image of the 5.9 nm SrTiO
3
and high-resolution TEM image of a single
SrTiO
3
nanocrystal shown at the inset.

37

synthesized on the gram scale to yield nanocrystals in near quantitative yields. The
overall compositions of the resulting solid-solution Ba
x
Sr
1-x
TiO
3
nanocrystals were
determined by elemental analysis, using inductively coupled plasma-optical emission
spectrometry (ICP-OES). Slight variations in the expected compositions compared to
actual compositions are likely as result of the different reaction rates of the two
bimetallic alkoxides. The synthesis method yields crystalline particles at neutral pH,
ambient pressure, and low temperatures; however, the reaction temperature of the
vapor diffusion sol-gel method had to be tuned depending on composition in order to
make solid-solution Ba
x
Sr
1-x
TiO
3
nanocrystals. For strontium-rich compositions
where 0 ≤ x ≤0.5, the reaction was performed at 80
º
C, and for barium-rich
compositions where 0.5 < x ≤ 1.0, the reactions were performed at 20
º
C. Performing
the reaction at higher temperatures for solid-solution compositions where x > 0.5
resulted in amorphous and ill-defined material in the same way that performing the
reaction at lower temperatures where x ≤ 0.5 resulted in amorphous and ill-defined
material. Thus, the temperature required for solid solution formation is governed by
the temperature needed to synthesize the pure ternary perovskite (i.e., BaTiO
3
or
SrTiO
3
) nanocrystals closest to its composition.
Powder X-ray diffraction (XRD) analysis of the SrTiO
3
nanocrystals revealed that
they are phase pure and are in the cubic perovskite phase, with a measured lattice
constant of a = 3.92 ± 0.02 Å, which is consistent with literature values (JCPDS no.
74-1296) (Figure 2.1a). The SrTiO
3
nanocrystal size was calculated to be ca. 6.4 nm
by Scherrer analysis of the XRD pattern, which is in close agreement with
38

measurements obtained from transmission electron microscopy (TEM) analysis (mean
diameter = 5.9 ± 1.2 nm) (Figure 2.1b). Lattice parameters for the Ba
x
Sr
1-x
TiO
3
(0 ≤ x
≤ 1) nanocrystals were calculated from the (110) reflection (100% intensity
reflection) relative to an internal silicon standard. As larger Ba
2+
cations are

Figure 2.2. (a) Powder X-ray diffraction pattern of the (110) reflection for the Ba
x
Sr
1-x
TiO
3

nanocrystals (0 ≤ x ≤ 1) referenced to the (111) reflection of an internal silicon standard (*). (b)
Linear dependence of the lattice parameter as a function of composition for the Ba
x
Sr
1-x
TiO
3

nanocrystals.

Figure 2.3. Low-resolution TEM images of (a) Ba
0.14
Sr
0.86
TiO
3
, (b) Ba
0.27
Sr
0.73
TiO
3
, (c)
Ba
0.33
Sr
0.67
TiO
3
, (d) Ba
0.43
Sr
0.57
TiO
3
, (e) Ba
0.48
Sr
0.52
TiO
3
, (f) Ba
0.63
Sr
0.37
TiO
3
, (g) Ba
0.69
Sr
0.31
TiO
3
,
and (h) Ba
0.77
Sr
0.23
TiO
3
nanocrystals.
39

introduced into the solid solution, the unit cell undergoes a systematic increase in
lattice parameter, as demonstrated by the shift of the (110) reflection to lower 2θ
values (Figure 2.2a). This gradual shift indicates a homogeneous distribution of
cations, forming a single crystalline solid solution as opposed to a polycrystalline
sample containing domains of both BaTiO
3
and SrTiO
3
. Furthermore, the solid
solution adheres to Vegard’s law, as evidenced by the linear dependence of the lattice
parameter in relation to the nanocrystal composition (Figure 2.2b).
18,19,23
Scherrer
analysis of the XRD pattern for the Ba
0.69
Sr
0.31
TiO
3
nanocrystals gives an estimated
crystallite size of 11.2 nm, which correlates well with TEM analysis (mean diameter =
11.6 ± 2.1 nm) (Figure 2.3g). Selected area electron diffraction analysis of the
Ba
0.69
Sr
0.31
TiO
3
nanocrystals indicates they are also crystalline and phase pure on a
more local scale (Figure 2.4). A high-resolution TEM image of an individual
Ba
0.69
Sr
0.31
TiO
3
particle with the (110) lattice planes displayed (d = 0.28 nm) suggests
the particles are single crystalline (Figure 2.4b). In general, the spacing between
lattice fringes measured for individual nanocrystals by high-resolution TEM were

Figure 2.4. (a) Selected area electron diffraction pattern for an ensemble of Ba
0.69
Sr
0.31
TiO
3

nanocrystals and (b) intensity line profile for lattice fringes of a single Ba
0.69
Sr
0.31
TiO
3

nanocrystal (inset).

40

homogeneous, again suggesting solid solution formation (Figure 2.4b inset). The
nanocrystals were on average relatively similar in size (within standard deviation)
through the entire composition range, with mean diameters ranging from 5.9-12.6 nm
(Table 2.1).
2.3.2 Dielectric Properties
The relative permittivity that bulk Ba
x
Sr
1-x
TiO
3
ceramics exhibit is known to vary
non-monotonically with composition, and a maximized relative permittivity is
observed when x ≈ 0.7 where local structural distortions are though to occur.
4,5,30
For
bulk Ba
x
Sr
1-x
TiO
3
ceramics, the non-monotonic variation in dielectric constant with
respect to composition results from the Ba
2+
cations’ deviation from Vegard’s law on
the local scale coupled with displacement of the Ti
4+
cations.
15
The substitution of
smaller Sr
2+
cations in place of Ba
2+
in the perovskite lattice results in local Ba
2+
-Ba
2+

interatomic spacings that do not decrease in a linear fashion when x ≥ 0.7. As more
Table 2.1. Mean sizes for Ba
x
Sr
1-x
TiO
3
nanocrystals.

Composition Mean diameter by TEM (nm) Size by Scherrer analysis (nm)
SrTiO
3
5.9 ± 1.2 6.4
Ba
0.14
Sr
0.86
TiO
3
8.8 ± 1.4 8.4
Ba
0.27
Sr
0.73
TiO
3
10.0 ± 1.3 11.0
Ba
0.33
Sr
0.67
TiO
3
9.5 ± 1.4 5.8
Ba
0.43
Sr
0.57
TiO
3
11.8 ± 1.5 10.4
Ba
0.48
Sr
0.52
TiO
3
9.9 ± 1.2 10.3
Ba
0.63
Sr
0.37
TiO
3
10.0 ± 1.5 11.8
Ba
0.69
Sr
0.31
TiO
3
11.6 ± 2.1 11.2
Ba
0.77
Sr
0.23
TiO
3
12.6 ± 1.8 12.7
BaTiO
3
5.9 ± 0.1 8.0

41

Sr
2+
cations are substituted into the perovskite lattice, a greater displacement of the
octahedrally coordinated Ti
4+
cations occurs, which leads to the higher relative
permittivity.
15
When x < 0.7, the local Ba
2+
-Ba
2+
interatomic spacing decreases
precipitously, leading to a constricted unit cell and a drastic decrease in the dielectric
constant.
Since the dielectric properties of perovskites are known to be size-dependent,
17-23,27
it
was of interest to study how the relative permittivity of small Ba
x
Sr
1-x
TiO
3

nanocrystals varies with composition. Previous studies into the dielectric properties
of nanocrystalline Ba
x
Sr
1-x
TiO
3
have shown a strong compositional dependence on
relative permittivity; however, these materials were sintered to high temperatures
prior to characterization of the dielectric properties.
18,23,31,32
Sintering has an effect
on the grain size – increasing size from tens of nanometers to potentially hundreds of
nanometers or larger via grain growth at high temperatures. To date, there has been
no direct study of the compositional dependent dielectric properties of small (sub-30
nm), unsintered Ba
x
Sr
1-x
TiO
3
nanocrystals. Although the particle size varies slightly
for the different compositions, the nanocrystals are all similarly sized within standard
deviation, and such small size differences should not have a dramatic effect on
relative permittivity.
The ensemble dielectric properties of the nanocrystals were studied by pressing a
mixture of the nanocrystals with 0.5 wt% polyvinyl alcohol (PVA) into cylindrical
pellets and coating the opposing faces with silver electrodes to form a simple parallel
42

plate capacitor. High temperature annealing was avoided to prevent particle sintering
and to maintain the original nanocrystal size, as confirmed by Scherrer analysis. To
obtain a more accurate measurement of the dielectric constant, Bruggeman’s effective
medium model was used;
33
0 = Σ ν[(ε
i
– ε
eff
)/(ε
i
+ 2ε
eff
)], where ν is the volume
fraction of the inclusion, ε
eff
is the dielectric constant of the composite pellet, and ε
i
is
the dielectric constant of the Ba
x
Sr
1-x
TiO
3
, PVA, and air inclusions. The dielectric
constants of polyvinyl alcohol and air were estimated as εʹ = 1.95 and 1.00,
respectively. To determine the volume fraction of air in each cylindrical pellet, the
pore volume of fractured pellets was measured by nitrogen porosimetry. The pore
volumes for the pellets of SrTiO
3
, Ba
0.69
Sr
0.31
TiO
3
, and BaTiO
3
were 3.1 x 10
-2
, 7.7 x
10
-2
and 6.7 x 10
-2
mL g
-1
, respectively (Figure 2.5).
This corresponds to the pellets having 13.5, 30.0, and 2.0% porosity for SrTiO
3
,
Ba
0.69
Sr
0.31
TiO
3
, and BaTiO
3
, respectively. Variation in the porosity of these samples
are likely a result of slight differences in particle size and morphology; however, the
degree of porosity for each of the samples is within the acceptable limits for
application of Bruggeman’s effective medium model.
33

Figure 2.5. Nitrogen adsorption (red)-desorption (black) isotherms for fractured pellets of (a)
SrTiO
3
, (b) Ba
0.69
Sr
0.31
TiO
3
, and (c) BaTiO
3
nanocrystals.
43

The dielectric properties of the Ba
x
Sr
1-x
TiO
3
nanocrystals were studied as a function of
solid-solution composition. In general, as a perovskite grain size decreases, the relative
permittivity also decreases.
17-23,27
As such, the measured dielectric constants (εʹ = 11.0–
341 at 1 kHz) for the Ba
x
Sr
1-x
TiO
3
nanocrystals are expectedly lower than the values for
bulk ceramic materials (εʹ = 5,000–20,000);
3
however, the same non-monotonic
dependence of relative permittivity on composition observed for bulk ceramics is also
observed for these small nanocrystals. It was observed that the dielectric constant of the
nanocrystals remained relatively unchanged as x approached 0.5; however, above x = 0.5,
there is an increase in dielectric constant of more than an order of magnitude as the
composition changes from x = 0.48 (εʹ = 18.6 at 1 kHz) to x = 0.69 (εʹ = 341 at 1 kHz),
and the dielectric constant significantly decreases when x > 0.7 (εʹ = 79.6 at 1 kHz when
x = 1.0) (Fig. 2.6). The general non-monotonic dependence of relative permittivity on
composition observed for bulk ceramics holds for these small, sub-15 nm nanocrystals.

Figure 2.6. Compositional dependence of the dielectric constant of Ba
x
Sr
1-x
TiO
3
nanocrystals
measured at 1 kHz. Measurements were taken at 20 °C.
44

Moreover, the relativity permittivity of the unsintered Ba
0.69
Sr
0.31
TiO
3
nanocrystals (εʹ =
341 at 1 kHz) prepared by the vapor diffusion sol-gel route is an order of magnitude
greater than the relative permittivity previously reported for unsintered 50-nm
Ba
0.69
Sr
0.31
TiO
3
nanocrystals (εʹ = 28 at 1 kHz).
32

Analysis of the SrTiO
3
nanocrystals revealed a dielectric constant (εʹ = 11.0 at 1 kHz)
that was lower than that of BaTiO
3
(εʹ = 79.6 at 1 kHz) and much lower than that of
the Ba
0.69
Sr
0.31
TiO
3
nanocrystals (εʹ = 340.7 at 1 kHz); however, the SrTiO
3

nanocrystals demonstrate a reduced frequency dependence of the dielectric constant
over the range of 1 kHz to 2 MHz—decreasing by 15% whereas the dielectric
constant for Ba
0.69
Sr
0.31
TiO
3
decreased by 41% within the same frequency range
(Figure 2.7a). The dielectric constant of the Ba
0.69
Sr
0.31
TiO
3
nanocrystals is more
than an order of magnitude greater than the SrTiO
3
nanocrystals at high frequencies.
The frequency dependence of the dielectric constant is a result of interfacial

Figure 2.7. (a) Dielectric constant of Ba
0.69
Sr
0.31
TiO
3
, BaTiO
3
, and SrTiO
3
nanocrystals as a
function of frequency. (b) Dielectric loss of Ba
0.69
Sr
0.31
TiO
3
, BaTiO
3
, and SrTiO
3
.
45

polarization at lower frequencies
17-23,34
has also been observed for BaTiO
3

nanocrystals for which the dielectric constant decreased 24% over the same frequency
range (Figure 2.7a).
35
This effect appears to be exacerbated in the case of the
Ba
0.69
Sr
0.31
TiO
3
nanocrystals and mitigated in the case of the SrTiO
3
nanocrystals.
The dielectric loss is significantly lower for SrTiO
3
nanocrystals (tan δ = 6.2 x 10
-2
at
1 kHz) than for Ba
0.69
Sr
0.31
TiO
3
(tan δ = 3.3 x 10
-1
at 1 kHz) or BaTiO
3
(tan δ = 1.3 x
10
-1
at 1 kHz) nanocrystals, and the SrTiO
3
nanocrystals also display a minimal
change in dielectric loss as the frequency was increased from 1 kHz to 2 MHz as
compared to the other nanocrystals (Figure 2.7b). The observed frequency stability of
the dielectric properties for the SrTiO
3
nanocrystals is consistent with reports from
the groups of Ganguli and Abraham who found a moderate change in dielectric
constant and dielectric loss over a range of frequencies and temperatures.
18,23
Despite
Ba
0.69
Sr
0.31
TiO
3
having highly frequency dependent dielectric loss, it is comparable to
the dielectric loss of BaTiO
3
nanocrystals (tan δ = 3.7 x 10
-2
for Ba
0.69
Sr
0.31
TiO
3

compared to tan δ = 2.8 x 10
-2
for BaTiO
3
at 2 MHz), but with a dielectric constant of
εʹ = 200 that is three times greater than that of the BaTiO
3
nanocrystals at high
frequency (i.e., at 2 MHz).
35

2.4 Conclusions
Small, well-defined Ba
x
Sr
1-x
TiO
3
nanocrystals have been made using an ultrabenign
synthesis method. The vapor diffusion sol-gel method provides the proper kinetically
controlled conditions to synthesize compositionally complex solid solution
46

nanocrystals at the lowest temperature reported heretofore. The ability to synthesize
these sub-15 nm nanocrystals over a wide composition range at low temperatures
allowed the effect of composition on dielectric properties to be studied for small,
unsintered nanocrystals. It was observed that the Ba
x
Sr
1-x
TiO
3
nanocrystals possess a
maximized dielectric constant at Ba
0.69
Sr
0.31
TiO
3
(εʹ = 341 at 1 kHz), which is more
than an order of magnitude greater than the dielectric constant of the pure SrTiO
3

nanocrystals, and higher than any other sub-30 nm Ba
x
Sr
1-x
TiO
3
nanocrystal
previously reported. The Ba
0.69
Sr
0.31
TiO
3
solid solution demonstrates a maximized
relatively permittivity as a result of local structural distortions, which has been
theoretically predicted and experimentally observed for bulk materials.
15,18
Thus, the
non-monotonic dependence of the dielectric constant on composition for bulk
perovskite solid solutions is also observed for sub-15 nm nanocrystals.
2.5 References
(1) Bhalla, S.; Guo, R.; Roy, R. Mat. Res. Innovat., 2000, 4, 3-26.
(2) Newnham, R. E.; Cross, L. E. MRS Bull., 2005, 30, 845-848.
(3) Principles and Applications of Ferroelectrics and Related Materials, ed. Lines,
M. E. and Glass, A. M. Clarendon Press, Oxford, 1977.

(4) Bunting, E. N.; Shelton, G. R.; Creamer, A. S. J. Am. Ceram. Soc., 1947, 30, 114-
125.

(5) Hilton, A. D.; Ricketts, B. W. J. Phys. D: Appl. Phys., 1996, 29, 1321-1325.
(6) Zhao, Z.; Buscaglia, V.; Viviani, M.: Buscaglia, T.; Mitoseriu, L; Testino, A.;
Nygren, M.; Johnsson, M.; Nanni, P. Phys. Rev. B, 2004, 70, 024107.

47

(7) Spanier, J. E.; Kolpak, A. M.; Urban, J. J.; Grinberg, I.; Ouyang, L.; Yun, W. S.;
Rappe, A. M.; Park, H. Nano Lett., 2006, 6, 735-739.

(8) Huang, L.; Chen, Z.; Wilson, J. D.; Banerjee, S.; Robinson, R. D.; Herman, I. P.;
Laibowitz, R.; O’Brien, S. J. Appl. Phys., 2006, 100, 034316.

(9) Smith, M. B.; Page, K.; Siegrist, T; Redmond, P. L.; Walter, E. C.; Sesahdri, R.;
Brus, L. E.; Steigerwald, M. L. J. Am. Chem. Soc. 2008, 130, 6955-6963.

(10) Caulton, K. G.; Hubert-Pfalzgraf, L. G. Chem. Rev., 1990, 90, 969-995.
(11) Chandler, C. D.; Roger, C.; Hampden-Smith, M. J. Chem. Rev., 1993, 93, 1205-
1241.

(12) Gao, Y.; Koumoto, K. Cryst. Growth Des., 2005, 5, 1983-2017.
(13) O’Brien, S.; Brus, L.; Murray, C. B. J. Am. Chem. Soc., 2001, 123, 12085-12086.
(14) Urban, J. J.; Yun, W. S.; Gu, Q.; Park, H. J. Am. Chem. Soc., 2002, 124, 1186-
1187.

(15) Mao, Y. B.; Banerjee, S.; Wong, S. S. Chem. Commun., 2003, 408-409.
(16) Wang, X.; Zhuang, J.; Peng, Q.; Li, Y. Nature 2005, 437, 121-124.
(17) Packia Selvam, I. P.; Kumar, V. Mater. Lett., 2002, 56, 1089-1092.
(18) Arya, P. R.; Jha, P.; Ganguli, A. K. J. Mater. Chem., 2003, 13, 415-423
(19) Hou, B.; Xu, Y.; Wu, D.; Sun, Y. Powder Technol., 2006, 170, 26-30.
(20) Chen, W. –P.; Zhu, Q. Mater. Lett., 2007, 61, 3378-3380.
(21) Jiquan, H.; Maochun, H.; Feilong, J.; Yongge, C. Mater. Lett., 2008, 62, 2304-
2306.

(22) Demirors, A. F.; Imhof, A. Chem. Mater., 2009, 21, 3002-3007.
(23) Pramanik, N. C.; Anisha, N.; Abraham, P. A.; Panicker, N. R. J. Alloy. Compd.,
2009, 476, 524-528.

(24) Niederberger, M.; Garnweitner, G.; Pinna, N.; Antonietti, M. J. Am. Chem. Soc.,
2004, 126, 9120-9126.

48

(25) Su, K.; Nuraje, N.; Yang, N. –L. Langmuir, 2007, 23, 11369-11372.
(26) Wei, X.; Xu, G.; Ren, Z.; Wang, Y.; Shen, G.; Han, G. J. Cryst. Growth 2008,
310, 4132-4137.

(27) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Small, 2008, 4, 2102-2106.
(28) Brutchey, R. L.; Morse, D. E. Angew. Chem., Int. Ed., 2006, 45, 6564-6566.
(29) Brutchey, R. L.; Cheng, G.; Gu, Q.; Morse, D. E. Adv. Mater., 2008, 20, 1029-
1033.

(30) Tanaka, H.; Tabata, H.; Ota, K.; Kawai, T. Phys. Rev. B, 1996, 53, 14112-14116.
(31) Zhang, L.; Zhong, W. –L.; Wang, C. –L.; Zhang, P. –L.; Wang, Y. –G. J. Phys.
D: Appl. Phys., 1999, 32, 546-551.

(32) Hornebecq, V.; Huber, C.; Maglione, M.; Antonietti, M.; Elissalde, C. Adv. Funct.
Mater., 2004, 14, 899-904.

(33) Tinga, W. R.; Voss, W. A.; Blossey, D. F. J. Appl. Phys., 1973, 44, 3897-3902.
(34) Hou, R. Z.; Ferreira, P.; Vilarinho, P. M. Chem. Mater., 2009, 21, 3536-3541.
(35) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Langmuir, 2010, 26, 5067-5071.

49

Chapter 3. Effect of Surface Modification on the Dielectric Properties
of BaTiO
3
Nanocrystals*

*Published in Langmuir 2010, 26, 5067.
3.1 Introduction
Bulk barium titanate (BaTiO
3
) exhibits ferroelectric behavior as a result of a distortion of
the ideal perovskite structure. In this noncentrosymmetric tetragonal phase (space group
P4mm), the Ti
4+
cation is displaced from the center of the unit cell relative to one of six
O
2–
anions, which results in a net dipole moment that can be oriented from one
crystallographic direction to another via an applied electric field.
1-3
The dipolar
properties of BaTiO
3
also make the material piezoelectric and pyroelectric, in addition to
having dielectric constants as high as   = 10,000.
4
Over the past several years, there
have been many successful attempts to synthesize BaTiO
3
nanocrystals, which display a
range of size-dependent ferroelectric, dielectric, and structural characteristics attributed to
increased surface energy.
5-11
The dynamics which govern the surface chemistry of
BaTiO
3
are largely dependent on the method of preparation and can result in the
formation of surface hydroxyls from sol-gel and hydrothermal techniques, carbonates
from chemisorbed CO
2
, or carbonaceous species from various precursor routes.
12-14

Controlled surface modification of BaTiO
3
is of interest for organic or aqueous solvent
dispersion, reduced particle aggregation, and tunable electronic properties.
15,16
For
example, Huber et al. observed a relationship between the dielectric loss (tan ) and the
thickness of a SiO
2
shell over Ba
1-x
Sr
x
TiO
3
. As the SiO
2
shell thickness increases, the
50

dielectric loss decreases and becomes more stable as a function of temperature when
compared to the unmodified material.
16
The insulating SiO
2
layer acts as a loss barrier
between the perovskite interfaces, thereby providing favorable loss characteristics.
The most common surface modifiers for metal oxide surfaces are carboxylates, silanes,
and phosphonates, each of which display various bonding modes depending on the
material's surface characteristics.
17
Carboxylates have been used as coordinating ligands
which chemisorb to the nanocrystal surface and aid in the formation of BaTiO
3
nanowires
and nanoparticles,
9,18
but they bind relatively weakly.
19
Alkoxy- and chlorosilanes are
surface modifiers that work via covalent condensation with hydroxyls on the nanocrystal
surface; however, silanes can self-condense to form ill-defined surface layers and
typically are very reactive towards water.
17
Phosphonates are known to have a strong
binding affinity for oxides, such as ITO, TiO
2
, Y
2
O
3
, and ZrO
2
, without having the
propensity to self-condense or react with water.
20-23
Perry and co-workers have used n-
octylphosphonates and various fluorinated arylphosphonates to aid in BaTiO
3
nanocrystal
dispersion within polymer nanocomposites.
19,24
Surface modification allows for particle
dispersion in organic solvents, which is ideal for solution processing of polymer
nanocomposites. In addition, the phosphonate-modified BaTiO
3
polymer
nanocomposites exhibit relatively high breakdown voltages at high particle loading. In
general, the dielectric strength of BaTiO
3
is notably poorer than most polymers, and is
one of the limiting factors in the energy storage performance of capacitors made from
composites incorporating these materials. While high particle loadings well above the
percolation threshold are required to achieve energy densities appreciably better than
51

pure polymer,
25
surface modifiers bearing medium- to long-chain or oligomeric pendent
moieties are believed to prevent direct particle contact and maintain a minimum
interparticle insulating layer regardless of volumetric loading.
Nanocomposites with high volume fractions of well-dispersed BaTiO
3
nanocrystals can
harness the dielectric, ferroelectric, and/or piezoelectric properties of BaTiO
3
while
compensating for its low breakdown voltage. To date, there has been no direct
investigation into the effect of organic surface modification on the dielectric properties of
BaTiO
3
nanocrystals as compared to their unmodified analogs. The differences in
dielectric properties of 6-nm BaTiO
3
nanocrystals both unmodified and modified with n-
hexylphosphonic acid (HPA) will be explored herein, with the results having important
implications in the development of polymer nanocomposites with surface-modified
nanocrystals.
3.2 Experimental Section
All manipulations were performed under nitrogen atmosphere using standard Schlenk
techniques. Air-free solvents were used throughout. Bimetallic
BaTi[OCH
2
CH(CH
3
)OCH
3
]
6
alkoxide (0.5 M in n-butanol/2-methoxypropanol (0.45/0.75
vol/vol); Gelest, Inc.) and n-hexylphosphonic acid (97%; Strem Chemicals, Inc.) were
used as received.

52

3.2.1 Synthesis of BaTiO
3
Nanocrystals
The BaTiO
3
nanocrystal synthesis was adapted from a literature procedure.
5
In short,
deionized water (20 mL) and a separate solution of bimetallic alkoxide (8.6 mL), diluted
to 0.43 M with n-butanol, were placed in an enclosed chamber under a nitrogen
atmosphere. The solution containing the bimetallic alkoxide was exposed to the water
vapor for 24 h. After 24 h, the water source was removed and the particles were allowed
to age under nitrogen at 37 ˚C for 24 h. The resulting off-white gel was collected and
rinsed in ethanol (3  10 mL) and dried in vacuo (25 ˚C, 0.05 mm Hg) for 24 h.
3.2.2 Surface Functionalization with n-Hexylphosphonic Acid
The as-prepared BaTiO
3
nanocrystals (414 mg) were suspended in 53.0 mL of dry
toluene and 3.55 mL of a 0.2 M solution of HPA in absolute ethanol was rapidly added.
The suspension was sonicated for 3 min and then heated to 90 ˚C for 24 h with stirring.
The product was isolated via centrifugation (6000 rpm for 20 min), washed with ethanol
(10 mL), toluene (10 mL), again with ethanol (10 mL), before being dried in vacuo (25
˚C, 0.05 mm Hg) for 24 h.
3.2.3 Material Characterization
Transmission electron microscopy (TEM) analysis was performed on a JEOL JEM-2100
microscope at an operating voltage of 120 kV, which was equipped with a Gatan Orius
CCD camera. Samples for TEM analysis were deposited from a suspension in methanol
for unmodified and in toluene for surface-modified nanocrystals on ultrathin carbon film
53

on a holey carbon support copper grids (Ted Pella, Inc.). Powder X-ray diffraction
(XRD) was performed on a Rigaku Ultima IV diffractometer using a Cu Kα (λ = 1.54 Å)
radiation source. BET measurements were performed on a Nova 2200e surface area and
pore size analyzer (Quantachrome Instruments, Inc.). Thermogravimetric Analysis
(TGA) was performed on a Shimadzu model TGA-50 at a ramp rate of 10
ᵒ
C/min to 1000
ᵒ
C in a nitrogen atmosphere. Phosphorus elemental analysis was performed at the
Microanalysis Laboratory at the University of Illinois at Urbana Champaign. FT-IR
spectra were collected on a Perkin-Elmer Spectrum 2000 at a scanning interval of 0.5 cm
-
1
with a resolution of 1 cm
-1
under flowing nitrogen. X-ray photoelectron spectroscopy
(XPS) was performed on a Surface Science M-Probe Spectrometer Model ESCA2703
using a monochromated Al anode. Solid-state
31
P NMR spectra were collected on a
Bruker DSX-500, operating at 202 MHz using a 4 mm CPMAS NMR probe. The sample
was spun at 12 kHz using dry air. A total of 8 scans were collected for HPA and 48 scans
for HPA-BaTiO
3
using a delay time of 100 s for both samples.
3.2.4 Dielectric Characterization
For ensemble dielectric studies, a 13-mm diameter pellet was prepared by grinding 200
mg of nanocrystals with 1 mL of an aqueous solution of polyvinylalcohol (PVA; 1
mg/mL) to give a 0.5 wt% loading of polymer. The slurry was allowed to dry under a
nitrogen atmosphere. The dried material was then reground and pressed with an applied
load of 10 metric tons in vacuo (10 mm Hg). The pellet was annealed under nitrogen at
150
ᵒ
C for 3 h. Colloidal silver paint (Ted Pella, Inc) was applied to both faces before
54

annealing the pellet at 100
ᵒ
C for 1 h under nitrogen. Capacitance and loss tangents were
measured on an Agilent 4294A impedance analyzer with a frequency sweep of 1 kHz to 2
MHz. Temperature studies were performed on a model 100 Integra Series Xtreme
hot/cold plate with liquid nitrogen coolant under a nitrogen atmosphere at a frequency of
2 kHz using a GW-Instek model LCR-816 capacitance meter.
3.3 Results and Discussion
3.3.1 Surface Modification of BaTiO
3
Nanocrystals
Barium titanate nanocrystals were prepared as previously reported by a vapor diffusion
sol-gel method, whereby a bimetallic BaTi[OCH
2
CH(CH
3
)OCH
3
]
6
alkoxide solution
(0.43 M in n-BuOH/2-methoxypropanol) was allowed to slowly react with water vapor
over the course of 24 h at room temperature.
5
This method of nanocrystal synthesis was
used instead of other synthetic methods because it does not require coordinating ligands

Figure 3.1. (a) Low-resolution TEM image of 6 nm HPA-BaTiO
3
nanocrystals and high-
resolution TEM image of a single HPA-BaTiO
3
nanocrystal as inset. (b) X-ray diffraction pattern
of cubic HPA-BaTiO
3
.

55

or a high temperature annealing step for crystallization, both of which would affect the
surface chemistry of the nanocrystals (vide supra). The as-prepared BaTiO
3
nanocrystals
are 6 nm in diameter and appear to be single crystalline.
5
Powder XRD analysis revealed
the nanocrystals are in the cubic perovskite phase, with a measured lattice constant of a =
4.04 ± 0.05 Å consistent with literature values (JCPDS no. 75-0215).
The as-synthesized BaTiO
3
nanocrystals were surface modified by heating the BaTiO
3

nanocrystals (414 mg) in toluene with 3.55 mL of a 0.2 M solution of HPA in absolute
ethanol at 90 ˚C. HPA-modified BaTiO
3
(HPA-BaTiO
3
) is stable to repeated solvent
rinsing with toluene and ethanol, which suggests a strong binding interaction between the
HPA and the nanocrystal surface. TEM analysis of the surface-modified nanocrystals
showed that the nanocrystal size and morphology are unaffected by reaction with HPA
(Figure 3.1a). Moreover, powder XRD revealed that the nanocrystals remain as phase-
pure BaTiO
3
after surface modification (Figure 3.1b).

Figure 3.2. TGA curves of unmodified and HPA-BaTiO
3
nanocrystals.
56

Qualitatively, the nanocrystals become much more dispersible in organic solvents (e.g.,
toluene) after surface modification with HPA, with dispersions of HPA-BaTiO
3
in
toluene remaining stable for several hours. This is to be expected since HPA binding
presumably breaks up particle agglomeration and adds a nonpolar organic corona to the
nanocrystal surface. Thermogravimetric analysis (TGA) of the HPA-BaTiO
3
showed a
weight loss between 430-450 ˚C under flowing nitrogen, which is attributed to loss of

Figure 3.4. High resolution XPS spectrum of the P 2p region of HPA-BaTiO
3
.

Figure 3.3. Nitrogen adsorption-desorption isotherm for (a) unmodified and (b) HPA-BaTiO
3
nanocrystals.
57

surface-bound HPA (Figure 3.2). Concomitant weight loss between 430 and 600
°
C for
the unmodified nanocrystals comes from residual organic species, such as chemisorbed
alkoxide functionality. The degree of surface modification was determined by BET
surface area analysis and TGA and corroborated by phosphorus elemental analysis. The
surface area of the unmodified BaTiO
3
nanocrystals was determined to be 123 m
2
/g.
After correcting for weight loss associated with the unmodified BaTiO
3
nanocrystals, the
degree of surface modification was calculated to be 2.4 phosphonate groups/nm
2
for
surface-modified HPA-BaTiO
3
. This represents ca. 57% of a theoretical monolayer
assuming an area of ~0.24 nm
2
for each phosphonate group.
22

Moreover, BET analysis of the unmodified BaTiO
3
nanocrystals revealed a type H2
absorption hysteresis, which is indicative of a non-uniform arrangement of rigidly joined
particles (Figure 3.3). BET analysis of HPA-BaTiO
3
, on the other hand, revealed a type
H4 absorption hysteresis, which is indicative of a loosely coherent assemblage of

Figure 3.5. FT-IR spectra of HPA, HPA-BaTiO
3
, and BaTiO
3
nanocrystals.

58

particles, as may be expected for nanocrystals with a dense organic corona.
26
In order to more completely investigate binding of HPA to the nanocrystal surface, XPS,
FT-IR, and
31
P MAS NMR spectroscopies were used. High resolution XPS indicated the
HPA was bound to the BaTiO
3
nanocrystals, with observation of P 2p
3/2
and 2p
1/2
peaks
at binding energies of 133.8 and 129.5 eV, respectively (Figure 3.4). The observed P 2p
binding energies are consistent with surface-bound phosphonates on ITO.
23
FT-IR
spectroscopy was also used to verify surface modification with HPA (Figure 3.5). The
HPA-BaTiO
3
sample shows strong ν(C–H) bands between 2965 and 2865 cm
-1
not
observed in the unmodified nanocrystals. Moreover, ν
a
(PO
3
2–
) and v
s
(PO
3
2–
) bands at
1060 and 990 cm
-1
, respectively, are consistent with tridentate phosphonate binding to the
nanocrystal surface.
27
The ν
a
(PO
3
2–
) and ν
s
(PO
3
2–
) bands differ significantly from the
series of ν(P=O) and ν(P–O) bands present between 1110 and 950 cm
-1
for free HPA. In
addition, the ν(PO–H) band at 2335 cm
-1
for free HPA is not present in the HPA-BaTiO
3
,

Figure 3.6.
31
P MAS NMR spectra of HPA and HPA-BaTiO
3
.

59

providing further evidence that the HPA binds to the nanocrystal surface in the
deprotonated phosphonate form.
28

31
P MAS NMR spectroscopy was also used to
elucidate binding of HPA to the nanocrystal surface (Figure 3.6). The
31
P MAS NMR
spectrum of free HPA revealed a sharp resonance at δ = 37.2 ppm, which is consistent
with other alkylphosphonic acids that have been reported in the literature.
19,22,29,30

The
31
P MAS NMR spectrum of HPA-BaTiO
3
revealed a strong resonance at δ = 21.7
ppm that is shifted upfield from that of free HPA. The change in chemical shift (Δδ =
15.5 ppm) is indicative of strong, and fixed tridentate binding of the phosphonate to the
nanocrystal surface.
22,29,30
In addition, a weaker, broad resonance between δ = 27.0-24.2
ppm was observed for HPA-BaTiO
3
, which is assigned to mono- and bidentate
coordination of the phosphonate to the nanocrystal surface. The broad resonance
suggests fluxional, or transient, species potentially situated over a wide range of different
surface environments.
19,22,29,30
Thus, FT-IR and
31
P MAS NMR spectroscopic analysis
suggests that the majority of HPA binds to the nanocrystal in a homogeneous tridentate
fashion, with a fraction of mono- and bidentate coordination as well. This is consistent
with previous studies of arylphosponic acids on TiO
2
surfaces, where binding occurs
mainly via tridentate coordination.
31

3.3.2 Dielectric Studies
A powder consisting of dried BaTiO
3
or HPA-BaTiO
3
nanocrystals was mixed with an
organic binder (0.5 wt% poly(vinyl alcohol)), pressed under high pressure, and coated
with colloidal silver electrodes to form a cylindrical parallel plate capacitor. High
60

temperature annealing was avoided to protect the integrity of the organic corona in HPA-
BaTiO
3
and to avoid particle sintering that occurs at higher temperatures. To obtain a
more accurate measurement of the dielectric constant, Bruggeman’s effective medium
model was used;
32
0 = Σ ν[(ε
i
– ε
eff
)/(ε
i
+ 2ε
eff
)], where ν is the volume fraction of the
inclusion, ε
eff
is the dielectric constant of the composite pellet, and ε
i
is the dielectric
constant of the BaTiO
3
, poly(vinyl alcohol), HPA, and air inclusions. The dielectric
constants of poly(vinyl alcohol), HPA, and air were estimated to be ε´ = 1.95, 1.60, and
1.00, respectively. To determine the volume fraction of air in each cylindrical pellet, the
pore volume of fractured pellets was measured by nitrogen porosimetry. The pore
volumes for the pellets of unmodified BaTiO
3
and HPA-BaTiO
3
were 6.7 x 10
-2
and 5.2 x
10
-3
mL/g, respectively (Figure 3.3). This corresponds to the pellets having 2 and 28%
porosity for the modified and unmodified BaTiO
3
, respectively.

Figure 3.7. (a) Dielectric constant of unmodified BaTiO
3
and HPA-BaTiO
3
as a function of
frequency. (b) Dielectric loss of unmodified BaTiO
3
and HPA-BaTiO
3
as a function of frequency.
Measurements were taken at 25 °C.
61

The dielectric properties of the unmodified BaTiO
3
and HPA-BaTiO
3
were measured as a
function of temperature and frequency to discern the effects of surface modification on
the material’s dielectric properties. It was observed that the dielectric constant of HPA-
BaTiO
3
(ε´ = 24.9 at 1 kHz) is lower than that of the unmodified BaTiO
3
(ε´ = 79.6 at 1
kHz); however, the HPA-BaTiO
3
has far superior frequency stability over the range of 1
kHz–2 MHz (Figure 3.7a,b). The dielectric constant of the unmodified BaTiO
3
decreases
by 24% as the frequency is increased over this range, whereas the dielectric constant of
the HPA-BaTiO
3
decreases by only 5% over the same frequency range. The large
frequency dependence of unmodified BaTiO
3
is likely a result of interfacial polarization
at lower frequencies,
33
which is mitigated in the case of HPA-BaTiO
3
. The dielectric loss
of HPA-BaTiO
3
(tan δ = 2.4 x 10
-2
at 1 kHz) was also found to be significantly lower
than that for unmodified BaTiO
3
(tan δ = 1.3 x 10
-1
at 1 kHz). Moreover, the dielectric
loss of HPA-BaTiO
3
was relatively insensitive to frequency as opposed to the unmodified

Figure 3.8. (a) Dielectric constant of unmodified BaTiO
3
and HPA-BaTiO
3
as a function of
temperature. (b) Dielectric loss of unmodified BaTiO
3
and HPA-BaTiO
3
as a function of
temperature. Measurements were taken at a frequency of 2 kHz.

62

BaTiO
3
(Figure 3.7b). The lower dielectric loss can be attributed to the organic corona
surrounding the nanocrystals, which acts as an insulating loss barrier that reduces
inelastic scattering of conducting charge carriers.
16,34
The temperature dependence of the dielectric constant and dielectric loss for unmodified
BaTiO
3
and HPA-BaTiO
3
was also studied (Figure 3.8). The dielectric constant of HPA-
BaTiO
3
remained relatively unchanged over a broad temperature range of -15 to 135 ˚C
(ε´
max
/ ε´
min
= 1.22 at 2 kHz), whereas the dielectric constant of unmodified BaTiO
3

displayed a much stronger, direct temperature dependence over the same temperature
range (ε´
max
/ ε´
min
= 1.96 at 2 kHz). While an increase in the dielectric constant is
observed for the unmodified BaTiO
3
, it is significantly weakened and broadened in
comparison to the Curie transition for large grain BaTiO
3
ceramics. The suppression and
broadening of the dielectric constant is consistent with previous reports by the groups of
Buscaglia and Steigerwald for BaTiO
3
nanocrystals.
35,36
The lack of a significant
temperature dependence and suppression of the dielectric constant for HPA-BaTiO
3

suggest extrinsic dielectric effects (resulting from surface modification) affect the
dielectric behavior. Both unmodified BaTiO
3
and HPA-BaTiO
3
exhibited increased
dielectric loss with increasing temperature, which can be attributed to a higher density of
freely promoted charge carriers at higher temperatures;
34
however, HPA-BaTiO
3
has a
less pronounced response to dielectric loss over the measured temperature range.

63

3.4 Conclusions
Small BaTiO
3
nanocrystals were functionalized with HPA in order to study the effects of
the phosphonate on the dielectric properties of the material. The dielectric constants and
dielectric loss of the unmodified BaTiO
3
and HPA-BaTiO
3
were studied as a function of
both frequency and temperature. The dielectric constant of HPA-BaTiO
3
was found to be
lower, but substantially less sensitive to frequency, than that of the unmodified BaTiO
3
–
decreasing only 5% as compared to 24% for unmodified BaTiO
3
over the frequency
range of 1 kHz–2 MHz. The temperature dependence of the dielectric constant at 2 kHz
was also found to be less sensitive for HPA-BaTiO
3
when compared to unmodified
BaTiO
3
. When analyzing the dielectric loss, HPA-BaTiO
3
was much lower (tan δ = 2.4 x
10
-2
at 1 kHz) and less sensitive to frequency and temperature when compared to
unmodified BaTiO
3
(tan δ = 1.3 x 10
-1
at 1 kHz). Huang et al. recently reported a
dielectric constant of ε´ = 85-90 and a dielectric loss of tan δ = 3.0 x 10
-2
for 10-30 nm
BaTiO
3
nanocrystals that had been sintered to 600 ˚C.
37
The dielectric constant was
found to be relatively insensitive to frequency, with the dielectric constant decreasing ca.
6% over the frequency range of 1-100 kHz. Sintering may affect the surface chemistry of
the nanocrystals such that low dielectric loss and greater frequency stability is realized
for similarly sized nanocrystals to those studied here. Surface modification with
phosphonates was shown to give superior dielectric properties in terms of lower dielectric
loss and greater temperature and frequency stability. Thus, surface-modification of
BaTiO
3
nanocrystals should be viewed as not only a method of tailoring the surface
chemistry, but also improving the dielectric quality of the material.
64

3.5 References
(1) Lines, M. E.; Glass, A. M. Principles and Applications of Ferroelectrics and
Related Materials; Clarendon Press: Oxford, 1977.

(2) Newnham, R. E.; Cross, L. E. MRS Bull. 2005, 30, 845-848.
(3) Chandler, C. D.; Roger, C.; Hampden-Smith, M. J. Chem. Rev. 1993, 93, 1205-
1241.

(4) Merz, W. J. Phys. Rev. 1949, 75, 687-687.
(5) Brutchey, R. L.; Morse, D. E. Angew. Chem. Int. Ed. 2006, 45, 6564-6566.
(6) Brutchey, R. L.; Cheng, G.; Gu, Q.; Morse, D. E. Adv. Mater. 2008, 20, 1029-
1033.

(7) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Small 2008, 4, 2102-2106.
(8) Mao, Y.; Banerjee, S.; Wong, S. S. Chem. Commun. 2003, 3, 408-409.
(9) O’Brien, S.; Brus, L.; Murray, C. B. J. Am. Chem. Soc. 2001, 123, 12085-12086.
(10) Gaskins, B. C.; Lannutti, J. J. J. Mater. Res. 1996, 11, 1953-1959.
(11) Niederberger, M.; Pinna, N.; Polleux, J.; Antonietti, M. Angew. Chem. Int. Ed.
2004, 43, 2270-2273.

(12) Jiang, B.; Peng, J. L.; Bursill, L. A. Ferroelectrics 1998, 207, 445-463.
(13) Spanier, J. E.; Kolpak, A. M.; Urban, J. J.; Grinberg, I.; Ouyang, L.; Yun, W. S.;
Rappe, A. M.; Park, H. Nano Lett. 2006, 6, 735-739.

(14) Bhalla, A. S.; Guo, R.; Roy, R. Mat. Res. Innovat. 2000, 4, 3-26.
(15) Chu, L. W.; Prakash, K. N.; Tsai, M. T.; Lin, I. N. J. Eur. Ceram. Soc. 2008, 28,
1205-1212.

(16) Huber, C.; Elissalde, C.; Hornebecq, V.; Mornet, S.; Treguer-Delapierre, M.;
Weill, F.; Maglione, M. Ceram. Int. 2004, 30, 1241-1245.

(17) Neouze, M. A.; Schubert, U. Montash. Chem. 2008, 139, 183-195.
65

(18) Urban, J. J.; Yun, W.; Gu, Q.; Park, H. J. Am. Chem. Soc. 2002, 124, 1186-1187.
(19) Kim, P.; Jones, S. C.; Hotchkiss, P. J.; Haddock, J. N.; Kippelen, B.; Marder, S.
R.; Perry, J. W. Adv. Mater. 2007; 19, 1001-1005.

(20) Mutin, P. H.; Guerrero, G.; Vioux, A. J. Mater. Chem. 2005, 15, 3761-3768.
(21) Traina, C. A.; Schwartz, J. Langmuir 2007, 23, 9158-9161.
(22) Gao, W.; Dickinson, L.; Grozinger, C.; Morin, F. G.; Reven, L. Langmuir 1996,
12, 6429-6435.

(23) Brewer, S. H.; Brown, D. A.; Franzen, S. Langmuir 2002, 18, 6857-6865.
(24) Kim, P.; Doss, N. M.; Tillotson, J. P.; Hotchkiss, P. J.; Pan, M. J.; Marder, S. R.;
Li, J.; Calame, J. P.; Perry, J. W. ACS Nano 2009, 3, 2581-2592.

(25) An, L.; Boggs, S. A.; Calame, J. P. IEEE Electr. Insul. M. 2008, 24, 5-10.
(26) Sing, K. S. W.; Everett, D. H.; Haul, R. A. W.; Moscou, L.; Pierotti, R. A.;
Rouquerol, J.; Siemieniewska, T. Pure Appl. Chem. 1985, 57, 603-619.

(27) Paniagua, S. A.; Hotchkiss, P. J.; Jones, S. C.; Marder, S. R.; Mudalige, A.;
Marrikar, F. S.; Pemberton, J. E.; Armstrong, N. R. J. Phys. Chem. C 2008, 112,
7809-7817.

(28) Stuart, B. H. Infrared Spectroscopy: Fundamentals and Applications; John Wiley
& Sons, Ltd.: Sussex, 2004.

(29) Gao, W.; Reven, L. Langmuir 1995, 11, 1860-1863.
(30) Mutin, P. H.; Guerrero, G.; Vioux, A. J. Mater. Chem. 2005, 15, 3761-3768.
(31) Lafond, V.; Gervais, C.; Maquet, J.; Prochnow, D.; Babonneau, F.; Hubert Mutin,
P. Chem. Mater. 2003, 15, 4098-4103.

(32) Tinga, W. R.; Voss, W. A. G. J. Appl. Phys. 1973, 44, 3897-3902.
(33) Hou, R. Z.; Ferreira, P.; Vilarinho, P. M. Chem. Mater. 2009, 21, 3536-3541.
(34) Kao, K. C. Dielectric Phenomena in Solids; Academic Press: London, 2004.
(35) Zhao, Z.; Buscaglia, V.; Viviani, M.; Buscaglia, M. T.; Mitoseriu, L.; Testino, A.;
Nygren, M.; Johnsson, M.; Nanni, P. Phys. Rev. B 2004, 70, 024107.
66

(36) Smith, M. B.; Page, K.; Siegrest, T.; Redmond, P. L.; Walter, E. C.; Seshadri, R.;
Brus, L. E.; Steigerwald, M. L. J. Am. Chem. Soc. 2008, 130, 6955-6963.

(37) Huang, L. M.; Chen, Z. Y.; Wilson, J. D.; Banerjee, S.; Robinson, R. D.; Herman,
I. P.; Laibowitz, R.; O’Brien, S. J. Appl. Phys. 2006, 100, 034316.

67

Chapter 4. Improved Breakdown Strength and Energy Density in Thin
Film Polyimide Nanocomposites with Small Barium Strontium Titanate
Nanocrystal Fillers*

*Published in J. Phys. Chem. C. 2013, 117, 6958.
4.1 Introduction
Commercially available capacitors are unable to meet the exacting demands of high
voltage pulsed power energy storage systems requiring high energy density, low
dielectric loss, and high temperature stability. Recently, polymer nanocomposites have
received a great deal of attention because of their potential to achieve high energy
densities. The underlying motivation behind using these hybrid nanocomposites is that
the polymer component will contribute high breakdown strength while the
nanoparticulate filler (e.g., TiO
2
, BaTiO
3
, etc.) will provide high relative permittivity.
The maximum energy density for a capacitor is given by U
v
= ½ ε
r
ε
0
E
BD
2
J/cm
3
, where ɛ
r

is the relative permittivity of the nanocomposite, ε
0
is the permittivity of free space, and
E
BD
is the breakdown strength of the nanocomposite.
1
Given the quadratic dependence of
energy density on the breakdown strength, significant gains in maximum energy density
can be made by maintaining or improving the composite breakdown strength relative to
the neat polymer while simultaneously adding to the relative permittivity. Unfortunately,
adding only a few weight percent of nanoparticulate filler can catastrophically reduce the
breakdown strength of the nanocomposite before significant improvements in the relative
permittivity are realized; therefore, the energy density of nanocomposites tends to
decrease with nanoparticulate filler concentration.
2-4

68

Polymer/BaTiO
3
nanocomposites typically have utilized large (30–100 nm)
nanoparticulate fillers with ill-defined morphologies, often resulting in inhomogeneous
nanocomposite films caused by gross filler agglomeration. To mitigate filler
agglomeration, Marder, Perry, and co-workers surface-modified commercial BaTiO
3

nanocrystals (30–120 nm) with phosphonic acids before extensive ball milling with
poly(vinylidenefluoride-co-hexafluoropropylene). In their system, the breakdown
strength of the neat polymer fell with inclusion of 5 vol% filler before a sharp drop-off
between 10 and 20 vol% filler, which was proposed to be a result of percolation onset.
5

More recently, Almadhoun et al. hydroxylated the surface of commercial BaTiO
3

nanocrystals with hydrogen peroxide and noted an improved breakdown strength for
nanocomposites with poly(vinylidene fluoride) (PVDF) over those with untreated fillers,
but the breakdown strength relative to the neat polymer still decreased beginning with
only 5 vol% filler loading.
6
Although these PVDF-based nanocomposites possess high
energy densities (as a result of the high relative permittivity of the majority polymer
component), these systems demonstrate relatively poor dielectric frequency stability, they
have inherently low thermal stability, and they have relatively high dielectric loss.
7
As a
result, polyimides are being explored as an alternative polymer component for
nanocomposites because of their high thermal and chemical stability, low dielectric loss,
and comparably high breakdown strength.
8-12
Dang et al. prepared thick (10–70 μm) freestanding films of Kapton nanocomposites via
an in-situ solution processing technique in which unmodified, 100-nm BaTiO
3

nanocrystals were pre-mixed with monomers prior to casting and thermal imidization.

69

Despite an intimate interaction between the poly(amic acid) (PAA) oligomers and the
BaTiO
3
nanocrystal surface, the breakdown voltage of the nanocomposites fell nearly
40% at loadings as low as 5 vol%; although, low dielectric loss and excellent thermal
stability were reported.
4
Feng et al. also prepared Kapton nanocomposites with 100-nm
unmodified BaTiO
3
nanocrystals and observed similar trends in decreasing breakdown
strength of the nanocomposites at low filler loading.
13
To date, there has been little focus on utilizing extremely small, well-defined perovskite
nanocrystals in these nanocomposite systems. The addition of small nanocrystalline
fillers into polymer matrices provide a means to enhance the dielectric properties beyond
that of the matrix material.
14-16
Small nanocrystals can provide substantial increases in
interfacial surface area and can be utilized to produce thinner films.
17-20
Larger micron-
sized fillers, on the other hand, have lower active surface areas and result in thicker and
generally more rough films.
21,22
While the room temperature dielectric constant of
BaTiO
3
is known to decrease with nanocrystal size,
23-25
this effect can be mitigated
through compositional tuning of the nanocrystal. For example, we have previously
reported that the relative permittivities of sub-15 nm Ba
x
Sr
1–x
TiO
3
and BaZr
x
Ti
1–x
O
3

nanocrystals can be maximized by synthetically controlling the cation ratios in the
perovskite lattice.
26,27
In this study, we demonstrate that the breakdown strength and
energy density of a novel polyimide/perovskite nanocomposite can exceed those values
of the neat polymer up to the percolation threshold. Thermoset nanocomposite films on
an ITO substrate were prepared by an in-situ polymerization method whereby well-
defined, sub-10 nm nanocrystals of Ba
0.7
Sr
0.3
TiO
3
(BST) were premixed with 1,3-bis(4-

70

aminophenoxy)benzene (BAPB) and pyromellitic dianhydride (PMDA) monomers prior
to spin-coating and subsequent imidization. The resulting thin films of the PMDA-
BAPB/BST nanocomposite demonstrated significantly improved breakdown strengths
and calculated energy densities more than twice that of the neat polyimide up to the
percolation threshold. To the best of our knowledge this is the first report of enhanced
breakdown strength with filler concentration in perovskite-based polymer composites.
4.2 Experimental Section
4.2.1 General Procedures
All manipulations were performed under a nitrogen atmosphere using dry, air-free
solvents throughout. Polyvinylpyrrolidone (PVP) (MW = 55,000; Sigma Aldrich, Inc.),
n-butanol (99.8%; Sigma Aldrich, Inc.), barium titanium bimetallic alkoxide
(BaTi[OCH
2
CH(CH
3
)OCH
3
]
6
; 0.5 M in n-butanol/3-methoxypropanol; Gelest, Inc.),
strontium titanium bimetallic alkoxide (SrTi[OCH
2
CH(CH
3
)OCH
3
]
6
; 0.7 M in n-
butanol/3-methoxypropanol; Gelest, Inc.), pyromellitic dianhydride (min. 98%; TCI
America, Inc.), 1,3-Bis(4-aminophenoxy)benzene (min. 98%; TCI America, Inc.), and
dimethylformamide (DriSolv; EMD Chemicals, Inc.) were used as received.
4.2.2 Vapor Diffusion Sol-Gel Synthesis of Ba
0.7
Sr
0.3
TiO
3
(BST) Nanocrystals
The details of this synthetic procedure have previously been reported elsewhere.
26-29
In
short, deionized water (20 mL) and 10 mL of a separate 0.43 M solution of barium
titanaium alkoxide (6.00 mL), strontium titanium alkoxide (1.84 mL), and n-butanol

71

(2.13 mL) were placed in an enclosed chamber under a nitrogen atmosphere. The
alkoxide solution was exposed to water vapor for 24 h at 20 ˚C. After 24 h, the water
was removed and the gel was allowed to age at 20 ˚C for an additional 24 h. The
resulting off-white gel was collected and rinsed in dimethylformamide (DMF) (3  10
mL), and finally dispersed in DMF at a concentration of 150 mg/mL. The BST
nanocrystal suspension was bath sonicated for 1 h under static nitrogen and stored at -22
˚C. The suspension remained stable in DMF after more than 6 months with very minimal
nanocrystal precipitation.
4.2.3 Material Characterization
Transmission electron microscopy (TEM) analysis was performed on a JEOL JEM-2100
microscope at an operating voltage of 200 kV, which was equipped with a Gatan Orius
CCD camera. Samples for TEM analysis were prepared by drop casting a suspension of
the nanocrystals in methanol onto ultrathin carbon film supported on 400 mesh copper
grids (Ted Pella, Inc.). Scanning electron microscopy (SEM) analysis was performed on
a JEOL JSM-6610LV instrument in high-vacuum mode using an accelerating voltage of
10 kV. Energy dispersive X-ray (EDX) spectroscopy analysis was performed using an
EDAX Apollo silicon-drift (model JSM 6490) mounted on a JEOL JSM-6610 SEM with
an accelerating voltage of 10 kV. Samples for SEM and EDX Spectroscopy were
prepared by spin casting a PMDA-BAPB/BST solution onto ITO coated glass and
annealing in an identical fashion to the tested devices (vide infra) and were then sputtered
with a thin layer of gold. BET measurements were performed on a Nova 2200e surface

72

area and pore size analyzer (Quantachrome Instruments, Inc.). Powder X-Ray diffraction
(XRD) was performed on a Rigaku Ultima IV diffractometer using Cu Kα (λ = 1.54 Å)
radiation. For XRD samples of nanocrystalline powder, the powder was added directly
onto a zero diffraction single crystalline silicon substrate (MTI Corporation, Inc.). For
XRD samples of PMDA-BAPB/BST nanocomposites, an aliquot of the PAA/BST
solution was cast onto a glass slide and thermally treated (vide infra). Thermogravimetric
analysis (TGA) was performed on a TA instruments Q50 Thermogravimetric analyzer
under both nitrogen and air environments at a heating rate of 10 °C/min. To measure the
dielectric constant of BST, a 13 mm diameter pellet was prepared by grinding 200 mg of
BST nanocrystalline powder with 1 mL of a PVP solution in methanol (1 mg/mL). The
slurry was allowed to dry overnight in a nitrogen atmosphere. The dried powder was
pressed with an applied load of 6 metric tons in vacuo (10 mm Hg). After pressing, the
pellet was heated under nitrogen at 100 °C for 3 h. Colloidal silver paint (Ted Pella, Inc.)
was applied to both sides before heating the pellet at 100 °C for 1 h under nitrogen. The
capacitance was measured at a frequency of 1 kHz using a GW-Instek model LCR-816
capacitance meter. In order to determine the dielectric constant, Bruggeman’s effective
medium model was used;
30
0 = Σ ν[(ε
i
– ε
eff
)/(ε
i
+ 2ε
eff
)], where v is the volume fraction of
the inclusion, ε
eff
is the dielectric constant of the composite pellet, and ε
i
is the dielectric
constant of the BST, PVP, and air inclusions. The dielectric constants of PVP and air
were estimated to be ε = 3.0 and 1.0, respectively, allowing for ε
i
(BST) to be solved as an
unknown. While the as-prepared BST nanocrystals are known to contain ~5-6 wt%
organics by TGA, the dielectric contribution of these organic species was not factored

73

into the effective medium model since their exact identity is unknown. Elemental
analysis was performed via inductively coupled plasma-mass spectrometry (ICP-MS) by
Galbraith Laboratories, Knoxville Tennessee.
4.2.4 Device Preparation
PMDA-BAPB/BST nanocomposites were fabricated via in-situ polymerization. First, an
aliquot of 150 mg/mL solution of BST in DMF was measured out and diluted to 3.5 mL
total volume in DMF (with the amount of BST used dependent upon desired loading).
This solution was then ultrasonicated with a Sonics Vibra-Cell 750 W ultrasonic
processor at 0 °C and 35 % amplitude for 1 h and immediately filtered through a 1.0 μm
syringe filter. An aliquot of the final filtered solution was used to determine the final
concentration and determine particle loading. To the filtered solution, 288 mg (0.99
mmol) of BAPB was added by continuous high-shear mixing for 10 min followed by the
addition of 215 mg (0.99 mmol) of PMDA in four equal aliquots separated by 10 min
with continuous high-shear mixing. After the final addition of PMDA, the concentration
of the solution was adjusted to the desired viscosity by dilution with DMF and was high-
shear mixed for an additional hour under flowing nitrogen. The resulting PAA/BST
nanocrystal suspension was then spun-cast onto etched tin-doped indium oxide (ITO)
coated glass substrates (7-10 Ω resistance; Colorado Concept Coatings, LLC.).
Substrates were cleaned by bath sonication for 5 min each in micro-organic soap,
deionized water, acetone, 2-propanol, and again in deionized water. The substrates were
then dried under flowing nitrogen and treated in an ozone atmosphere for 20 min. Upon

74

casting, the films were sequentially heated in vacuo (0.05 mmHg) under the following
conditions: 90 min each at 150, 200, and 250 °C followed by heating for 12 h at 200 °C.
The films were heated at 300 °C for 2 h under flowing nitrogen and were allowed to
slowly cool to room temperature and stored in a dry, air-free environment. Film
thickness ranged between 2-3 μm as measured by profilometry (AmBios XP2 stylus
profilometer with 2.0 μm tip diameter; Ambios Technologies, Inc.).
4.2.5 Device Characterization
Parallel plate capacitors were fabricated by depositing an array of finger-like (16.45
mm
2
) top Al electrodes on the PMDA-BAPB/BST nanocomposite thin films. Aluminum
(250 nm thickness) was deposited through a shadow mask by using a custom thermal
evaporator at a deposition rate of 3 Å/s. Capacitance and loss tangents were measured on
an Agilent 4294A impedance analyzer with a frequency sweep of 1 kHz – 1 MHz. DC
negative polarity breakdown measurements were performed at a linear ramp rate of 250
V/s (Bertan 225, 20 kV DC power supply), controlled by a custom made bias box (Figure
4.1) and monitored by an oscilloscope (Tektronix TDS 2004C; Tektronix, Inc.).
4.3 Results and Discussion
4.3.1 Synthesis, Characterization, and Dielectric Properties of BST
BST nanocrystals were prepared by a vapor diffusion sol-gel technique whereby a
mixture of barium titanium and strontium titanium bimetallic alkoxides (i.e.,
ATi(OCH
2
CH(CH
3
)OCH
3
)
6
where A = Ba or Sr) were exposed to water vapor at room

75

temperature to yield crystalline BST nanoparticles.
26-29
The resulting nanocrystals
formed stable colloidal suspensions in DMF after sonication. Recovered nanocrystal
yields were quantitative; however, it is important to note that the nanocrystals were never
brought to dryness for practical purposes, as this caused irreversible agglomeration and
prevented filtration for device fabrication. The resulting BST nanocrystals were phase
pure and indexed to the Pm m cubic perovskite phase,
26
with a measured lattice constant
of a = 4.01 Å as confirmed by powder X-ray diffraction (XRD) (Figure 4.2). The

3

Figure 4.1. Schematic of high voltage testing station used for all breakdown measurements.

76

elemental composition of the BST nanocrystals was confirmed to be 36:14:50 Ba/Sr/Ti
by ICP-MS. Transmission electron microscopy (TEM) analysis of the as-prepared
nanocrystals revealed that the particles are well defined with an average diameter of 9.5 ±
1.4 nm (Figure 4.3a). High resolution TEM images of the BST nanocrystals clearly show
the {110} and {100} family of planes measuring d = 2.79 and 4.00 Å, respectively,
consistent with XRD results (Figure 4.3b,c). Ensemble dielectric measurements of the
resulting unsintered BST nanocrystals gave a relative permittivity of ɛ
r
= 315, which is

Figure 4.2. X-ray diffraction patterns of BST and PMDA-BAPB/BST nanocomposite thin films.

Figure 4.3. (a) Low resolution TEM image of BST nanocrystals, and (b,c) high-resolution TEM
images showing the {110} and {100} family of lattice planes.

77

greater than the relative permittivity previously reported for similarly sized BaTiO
3

nanocrystals (ɛ
r
= 50-80) prepared by the same synthesis route.
4.3.2 PMDA-BAPB/BST Nanocomposite Fabrication and Characterization
Dispersed BST nanocrystals were premixed with BAPB and PMDA monomers to form a
PAA/BST blend in DMF. The nanocomposites were fabricated by spin-casting the
PAA/BST blend onto ITO-coated glass substrates followed by slow thermal imidization
under vacuum up to 300 °C (vide supra) to produce 2-3 µm PMDA-BAPB/BST thin
films. Low molecular weight PAA oligomers formed during the polymerization have
been postulated to aid in dispersion by improving the oxide-polymer interfacial
interaction,
4
since carboxylic acids are known to bind strongly to BaTiO
3
nanocrystal
surfaces.
31,32
Additionally, the rapid drying achieved by spin-casting helped to
kinetically immobilize the nanocrystalline filler and prevent particle settling that would
occur through solution casting routes. The parent PMDA-BAPB polyimide was chosen
because the glass transition temperature (T
g
) of the polyimide was lower than Kapton

Figure 4.4. Digital photograph of all PMDA-BAPB/BST nanocomposites.

78

which reduced processing temperatures and provided a more flexible film, while still
maintaining high thermal stability (>500 °C), low loss (tan δ = 0.015), and a modest
dielectric constant ( ε
r
= 2.8).
33-36
Variable concentrations of BST suspensions were
added such that the resulting PMDA-BAPB/BST nanocomposites possessed nanocrystal
loadings of 0, 5, 10, 13, 15, and 18 vol%. The as-prepared nanocomposite thin films
were yellow in color (like the parent polyimide) and were optically transparent,
suggesting a fairly homogeneous dispersion of the nanocrystals within the polyimide
matrix (Figure 4.4). Scanning electron microscopy (SEM) images of the nanocomposite
surface showed relatively uniform film morphology, and energy dispersive X-ray (EDX)
spectroscopic mapping of the Ti K emission lines on the film surface corroborated a
homogeneous distribution of titanium cations within the nanocomposite (Figure 4.5).
XRD analyses of the nanocomposite thin films revealed the BST nanocrystals remained

Figure 4.5. SEM micrographs with Ti elemental mapping (shown as inset) for PMDA-
BAPB/BST nanocomposites at 5 (a), 10 (b), 13 (c), 15 (d), and 18 (e) vol% BST loading.

79

phase pure after processing and fabrication, with no indication of BaCO
3
or SrCO
3

formation (Figure 4.2).
The PMDA-BAPB/BST composite system was thermally robust, as confirmed by
thermogravimetric analysis (TGA). Addition of the BST nanoparticulate filler resulted in
a modest decrease in thermal stability relative to the neat PMDA-BAPB polyimide,
which possessed a decomposition onset (determined by the first derivative of the wt% as
a function of temperature) at 560 °C in air and 540 °C in nitrogen by TGA. The resulting
PMDA-BAPB/BST nanocomposites were all stable up to 450 °C in air and 500 °C in
nitrogen (Figure 4.6). The larger variation in thermal stability in air is a result of catalytic
decomposition of the polyimide in the presence of oxides.
9,37
FT-IR spectroscopy was used to verify complete imidization of the neat PMDA-BAPB
polyimide under the processing conditions (Figure 4.7). The appearance of strong
ν(C=O) in-phase and out-of-phase imide stretching modes at 1780 and 1731 cm
–1
,

Figure 4.6. TGA curves in air (a) and in nitrogen (b) for PMDA-BAPB/BST nanocomposites
at all loadings.

80

respectively, and ν(C-N-C) axial and transverse stretching and out-of-plane deformation
modes at 1375, 1169, and 725 cm
–1
, respectively, are indicative of imidization.
38,39

Furthermore, the IR bands associated with the PAA intermediate (e.g., ν(C=O) acid or
amide modes at 1720 and 1670 cm
-1
, respectively, or v(C-N-H) amide mode at 1545 cm
-
1
) were not observed for the neat polymer. As confirmation of the FT-IR data, the neat

Figure 4.7. FT-IR spectrum of the neat PMDA-BAPB polyimide.

Figure 4.8. Raman spectra of PMDA-BAPB/BST nanocomposites, offset for clarity. The
neat PMDA-BAPB film was processed and prepared in an analogous way to the
nanocomposites but without addition of BST.

81

polymer displayed prominent Raman bands at 1512 and 1604 cm
-1
corresponding to the
ν(C
6
H
4
) modes and a band at 1788 cm
-1
that is referenced to the ν(C=O) in-phase imide
mode, with no PAA present (Figure 4.8).
38-41
As the loading of BST increases within the
polymer matrix, a characteristic ν(C-N-H) amide band appears with increasing intensity
at 1560 cm
-1
as a result of incomplete imidization. This phenomenon has been previously
reported for polyimide/BaTiO
3
nanocomposites and attributed to the nanocrystals
interacting with PAA, reducing chain mobility, and consequently hindering imidization.
9

Despite incomplete imidization and the presence of amide moieties within the
nanocomposite, the small BST nanocrystals are able to provide improved device
performance (vide infra).
4.3.3 Dielectric Characterization
The dielectric constant and loss tangent of the nanocomposites were measured at a range
of frequencies from 1 kHz to 1 MHz. The dielectric constant of the PMDA-BAPB/BST
nanocomposite increased with increasing BST loading over all frequencies tested. For
example, the dielectric constant increased from ɛ
r
= 2.8 for the neat PMDA-BAPB
polyimide to ɛ
r
= 6.2 at 18 vol% BST loading (measured at 1 MHz, Figure 4.9a). The
measured dielectric constant of the PMDA-BAPB/BST nanocomposite as a function of
BST loading was compared against those values predicted by Bruggeman's effective
medium model, 0 = Σ ν[(ε
i
– ε
eff
)/( ε
i
+ 2ε
eff
)],
30
where v is the volume fraction of the
inclusion, ɛ
eff
is the dielectric constant of the nanocomposite, and ɛ
i
is the dielectric
constant of BST and PMDA-BAPB, which were measured to be 315 and 2.8,

82

respectively. As seen in Figure 5a, the experimentally measured data matches well with
those values predicted by the effective medium model within this range of BST loadings.
Moreover, the PMDA-BAPB/BST nanocomposites exhibited excellent dielectric stability
over these frequencies (Figure 4.10a,b), which is an improvement over P(VDF-
HFP)/BaTiO
3
nanocomposite systems.
5,18
In each of the BST nanocrystal loadings
tested, the dielectric loss of the nanocomposites remained below 0.04 over the frequency
range of 1 kHz to 1 MHz. Such values are consistent with other perovskite-loaded
polyimide composites,
4,13
and is markedly improved over PVDF systems which have
shown dielectric losses as high as 0.2 at 1 MHz.
5

To study the effects of the BST nanocrystals on the dielectric breakdown strength of the
PMDA-BAPB/BST nanocomposite, the thin films were subjected to a negative DC bias
at a ramp rate of 250 V/s until breakdown was indicated by an instantaneous increase in

Figure 4.9. (a) The Weibull breakdown at 63.2 % probability of failure (red, left axis) and
relative permittivity compared to Bruggeman’s effective medium model (blue, right axis) as a
function of BST loading. (b) Two-parameter Weibull plots of PMDA-BAPB/BST
nanocomposites where dashed line represents Weibull breakdown at 63.2 % probability of
failure. Measurements were taken at 25 °C.

83

current. Dielectric breakdown strength was calculated using a two-parameter Weibull
distribution function:
42
P = 1- exp[-(E/E
BD
)
ß
], where P is the cumulative probability of
breakdown, E is the experimentally measured breakdown strength, E
BD
is the cumulative
breakdown probability at 63.2%, and β is a shape parameter for 15-20 independent
measurements per volume fraction. The Weibull distributions of breakdown strengths are
given in Figure 4.9a and b. The breakdown strength of the PMDA-BAPB/BST
nanocomposites increases with BST loading up to 10 vol%, where it reaches a maximum
value of 296 V/µm. This represents a 24% increase over the breakdown strength of the
neat PMDA-BAPB polyimide (E
BD
= 238 V/µm). At BST nanocrystal volume fractions
above 10%, the breakdown strength falls slightly to become comparable to that of the
neat PMDA-BAPB before falling below that of the neat polymer between 15 and 18
vol% BST loading. The precipitous decrease in breakdown strength at these BST volume
fractions is likely the result of a continuous network of BST nanocrystals percolating
across the electrodes.
14,43
In the percolation regime, E
BD
is dominated by the
nanoparticulate filler and negatively impacts the high breakdown strength of the matrix
material.
43,44

While an enhancement in breakdown strength from nanoparticulate fillers below the
percolation threshold has been reported in several nanocomposite systems,
45-51
such an
effect has not been previously reported for any perovskite-based nanocomposites.
Nanocomposites are known to possess a high interfacial area because of the high surface
area of nanoparticulate fillers relative to larger fillers;
4
in turn, the interface between the
nanocrystals and polymer will be dominant even at low volume fractions.
52,53
This short-

84

range layer surrounding the nanoparticulate fillers is thought to allow charge dissipation
and improve the internal electric field distribution, which help to suppress significant
interfacial polarization that exists in the case of larger fillers.
54,55
The dispersed
nanoparticulate fillers also help decrease charge transport by acting as scattering
centers.
47,54,55
The mitigating effects of nanoparticulate fillers can help improve the
breakdown strength of the resulting nanocomposites. Here, the excellent dispersion of
sub-10 nm BST nanocrystals within the PMDA-BAPB matrix has allowed for an increase
in breakdown strength relative to the neat polymer up to 10 vol% loading. This may be
rationalized by the strong interaction between the BST nanocrystals and the polymer
matrix (vide supra), which likely reduces the localized space charge accumulation.
3,54,56

As a result of favorable effects inherent to the nanoparticulate nature of the BST filler,
the calculated energy density of the PMDA-BAPB/BST increased from 1.4 J/cm
3
for the
pure polyimide to 2.9 J/cm
3
at 10 vol% BST, which represents an increase of 107% (see
Supporting Information, Figure S7). Such improvements have never been realized for
perovskite nanocrystals, where precipitous drops in breakdown strength with marginal
increase in permittivity result in lowered energy density at low filler volume fraction.
These results suggest that perovskite nanocrystals in the size regime of ~10 nm may be
needed to reap these benefits, since previously reported polymer/perovskite
nanocomposites with fillers as small as 30 nm showed diminished breakdown strength at
low volume fractions;
2,4,5,13
however, it is likely the interplay between filler size and
dispersion within the polymer matrix dominates these interfacial effects.

85

4.4 Conclusions
Nanocrystals of BST have been successfully blended into a PMDA-BAPB polyimide
system to understand the effects of loading small perovskite fillers on the dielectric
properties of the nanocomposite. The relative permittivity of the resulting PMDA-
BAPB/BST nanocomposites continually increased from ɛ
r
= 2.8 for the neat PMDA-
BAPB polyimide to ɛ
r
= 6.2 at 18 vol% BST, while tan δ <0.04 was maintained for all
compositions up to 1 MHz. The breakdown strength of the PMDA-BAPB/BST
nanocomposites increased with BST loadings up to 10 vol%, where a maximum
breakdown strength of 296 V/µm was reached. As a result, the calculated energy density
improved by more than twice that of the neat PMDA-BAPB polyimide at 10 vol% BST.
The simultaneous increase in relative permittivity with a decrease in dielectric loss
observed for the 10 vol% nanocomposite, combined with the increased breakdown
strength, make this composition particularly significant for practical applications. Such

Figure 4.10. (a) Relative permittivity of PMDA-BAPB/BST devices as a function of frequency.
(b) Dielectric loss of PMDA-BAPB/BST devices as a function of frequency. Measurements
were taken at 25 °C.

86

improvements have never been realized for perovskite-based nanocomposites and are
likely a result of enhanced interfacial interactions that mitigate space charge build-up and
suppress interfacial polarization, and/or scattering effects. The use of extremely small
(sub-10 nm), well-dispersed nanocrystals is beneficial in realizing the potential of
nanocomposites for capacitive energy storage solutions.
4.5 References
(1) Barber, P.; Balasubramanian, S.; Anguchamy, Y.; Gong, S.; Wibowo, A.; Gao,
H.; Ploehn, H. J.; Loye, H. C. Materials, 2009, 2, 1697-1733.

(2) Khalil, M. S. IEEE T. Dielect. El. In. 2000, 7, 261-268.

(3) Ma, D.; Hugener, T. A.; Siegel, R. W.; Christerson, A.; Martensson, E.; Onneby,
C.; Schadler, L. S. Nanotechnology 2005, 16, 724-731.

(4) Dang, Z. M.; Lin, Y. Q.; Xu, H. P.; Shi, C. Y.; Li, S. T.; Bai, J. Adv. Funct.
Mater. 2008, 18, 1509-1517.

(5) Kim, P.; Doss, N. M.; Tillotson, J. P.; Hotchkiss, P. J.; Pan, M. J.; Marder, S. R.;
Li, J.; Calame, J. P.; Perry, J. W. ACS Nano 2009, 3, 2581-2592.

(6) Almadhoun, M. N.; Bhansali, U. S.; Alshareef, H. N. J. Mater. Chem. 2012, 22,
11196-11200.

(7) Zhou, T; Zha, J. W.; Cui, R. Y.; Fan, B. H.; Yuan, J. K.; Dang, Z. M. ACS Appl.
Mater. Interfaces 2011, 3, 2184-2188.

(8) Fragiadakis, D.; Logakis, E.; Pissis, P.; Kramarenko, Y. V.; Shantalii, T. A.;
Karpova, I. L.; Dragan, K. S.; Privalko, E. G.; Usenko, A. A.; Privalko, V. P. J.
Phys. Conf. Ser. 2005, 10, 139-142.

(9) Lin, B.; Liu, H.; Zhang, S.; Yuan, C. J. Solid State Chem. 2004, 177, 3849-3852.

(10) Chen, D.; Zhu, H.; Liu, T. ACS Appl. Mater. Interfaces 2010, 2, 3702-3708

(11) Venkatasubramanian, N.; Wiacek, K. J.; Fries-Carr, S.; Fossum, E.; Dang, T. D.
High-Temperature Polymer Dielectrics for Capacitive Energy-Storage

87

Applications. In Polyimides and Other High Temperature Polymers: Synthesis,
Characterization and Applications, Volume 4; Mittal, K. L., Ed.; Brill: Boston,
Massachusetts, 2010; pp 393-406.

(12) Murugaraj, P.; Mainwaring, D.; Mora-Huertas, N. J. Appl. Phys. 2005, 98, 1-6.

(13) Feng, Y.; Yin, J.; Chen, M.; Liu, X.; Li, G. 6
th
International Forum on Strategic
Technogology; Heilongjian, China; August 22-24, 2011; pp 226-229.

(14) Roy, M.; Nelson, J. K.; MacCrone, R. K.; Schadler, L. S.; Reed, C. W.; Keefe, R.;
Zenger, W. IEEE T. Dielect. El. In. 2005, 12, 629-643.

(15) Nanophase and Nanocomposite Materials II; Komarneni, S.; Parker, J. C.;
Wollenberger, H. J., Eds.; MRS Symposium Proceedings 457; Materials Research
Society: Pittsburg, PA, 1997.

(16) Nilofar, A. Nanocomposites Preparation: Prospectives and Limitations. In
Nanocomposites Preparation, Properties, and Performance; Mancini, L. H.;
Esposito, C. L., Eds. Nova Science: New York, 2008; pp 241-251.

(17) Tchoul, M. N.; Fillery, S. P.; Koerner, H.; Drummy, L. F.; Oyerokun, F. T.;
Mirau, P. A.; Durstock, M. F.; Vaia, R. A. Chem. Mater. 2010, 22, 1749-1759.

(18) Tang, H.; Ma, Z.; Zhong, J.; Yang, J.; Zhao, R.; Liu, X. Colloid Surface A 2011,
384, 311-317.

(19) Xie, S. H.; Zhu, B. K.; Wei, X. Z.; Xu, Z. K.; Xu, Y. Y. Compos. Part A.-Appl. S.
2005, 36, 1152-1157.

(20) Xie, L.; Huang, X.; Wu, C.; Jiang, P. J. Mater. Chem. 2011, 21, 5897-5906.

(21) Althues, H.; Henle, J.; Kaskel, S. Chem. Soc. Rev. 2007, 36, 1454-1465.

(22) An, L. A.; Boggs, S. A.; Calame, J. IEEE Int. Sym. Elec. In. 2006, 24, 466-469.

(23) Huang, L.; Chen, Z.; Wilson, J. D.; Banerjee, S.; Robinson, R. D.; Herman, I. P.;
Laibowitz, R.; O’Brien, S. J. Appl. Phys. 2006, 100, 034316.

(24) Buscaglia, V.; Buscaglia, M. T.; Vivani, M.; Mitoseriu, L.; Nanni, P.; Trefiletti,
V.; Piaggio, P.; Gregora, I.; Ostapchuk, T.; Pokorny, J. et al. J. Eur. Ceram. Soc.
2006, 26, 2889-2898.

(25) Zhao, Z.; Buscaglia, V.; Vivaiani, M.; Buscaglia, M. T.; Mitoseriu, L.; Testino,
A.; Nygren M.; Johnsson, M.; Nanni, P. Phys. Rev. B. 2004, 70, 024107.

88

(26) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. J. Mater. Chem. 2010, 20, 5074-
5079.

(27) Rabuffetti, F. A.; Brutchey, R. L. Chem. Commun. 2012, 48, 1437-1439.

(28) Rabuffetti, F. A.; Lee, J. S.; Brutchey, R. L. Chem. Mater. 2012, 24, 3114-3116.

(29) Brutchey, R. L.; Morse, D. E. Chem. Int. Ed. 2006, 45, 6564-6566.

(30) Tinga, W. R.; Voss, W. A. G.; Blossey, D. F. J. Appl. Phys. 1973, 44, 3897-3902.

(31) Urban, J. J.; Yun, W.; Gu, Q.; Park, H. J. Am. Chem. Soc. 2002, 124, 1186-1187.

(32) O’Brien, S.; Brus, L.; Murray, C. B. J. Am. Chem. Soc. 2001, 123, 12085-12086.

(33) Hsiao, S. H.; Chen, Y. J. Eur. Polym. J. 2002, 38, 815-828.

(34) Ishii, J.; Takata, A.; Oami, Y.; Yokota, R.; Vladimirov, L.; Hasegawa, M. Eur.
Polym. J. 2010, 46, 681-693.

(35) Tamai, S.; Yamaguchi, A.; Ohta, M. Polymer 1996, 37, 3683-3692.

(36) Chao, M.; Kou, K. C.; Wang, Z. C.; Guang, L. W.; Zhang, D. N.; Zang, J. Q. Adv.
Mat. Res. 2011, 199, 13-18.

(37) Tong, Y.; Li, Y.; Xie, F.; Ding, M. Polym. Int. 2000, 49, 1543-1547.

(38) Young, J. T.; Tsai, W. H.; Boerio F. J. Macromolecules 1992, 25, 887-894.

(39) Ishida, H.; Wellinghoff, S. T.; Baer, E.; Koenig, J. L. Macromolecules 1980, 13,
826-834.

(40) Konieczny, J. M.; Wunder, S. L. Macromolecules 1996, 29, 7613-7615.

(41) Ge, J. J.; Xue, G.; Li, F.; McCreight, K. W.; Want, S. Y.; Harris, F. W.; Cheng, S.
Z. D.; Zhuang, X.; Hong, S. C.; Shen, Y. R. Macromol. Rapid Commun. 1998, 19,
619-623.

(42) Dissado, L. A.; Fothergill, J. C. Electrical Degradation and Breakdown in
Polymers; Stevens, G. C., Ed.; IEEE Materials and Devices 9; Peter Peregrinus
Ltd.: London, United Kingdom, 1992; pp 323-329.

89

(43) Coppard, R. W.; Bowman, J.; Dissado, L. A.; Rowland, S. M.; Rakowski, R. T. J.
Phys. D: Appl. Phys. 1990, 23, 1554-1561.

(44) Momen, G.; Farzaneh, M. Rev. Adv. Mater. Sci. 2011, 27, 1-13.

(45) Tuncer, E.; Sauers, I.; James, D. R.; Ellis, A. R.; Paranthaman, M. P.; Goyal, A.;
More, K. L. Nanotechnology 2007, 18, 325704-325708.

(46) Tuncer, E.; Rondinone, A. J.; Woodward, J.; Sauers, I.; James, D. R.; Ellis, A. R.
Appl. Phys. A. 2009, 94, 843-852.

(47) Fillery, S. P.; Koerner, H.; Drummy, L.; Dunkerley, E.; Durstock, M. F.; Schmidt,
D. F.; Vaia, R. A. ACS Appl. Mater. Inter. 2012, 4, 1388-1396.

(48) Ma, D.; Siegel, R. W.; Hong, J. I.; Schadler, L. S.; Martensson, E.; Onneby, C. J.
Mater. Res. 2004, 19, 857-863.

(49) Tomer, V.; Polizos, G.; Manias, E.; Randall, C. A. J. Appl. Phys. 2010, 108,
074116.

(50) Brandstetter, S. S.; Drummy, L. F.; Horwath, J. C.; Schweickart, D. L.; Vaia, R.
V. IEEE T. Dielect. El. In. 2008, 15, 287-290.

(51) Imai, T.; Sawa, F.; Nakano, T.; Ozaki, T.; Shimizu, T.; Kozako, M.; Tanaka, T.
IEEE T. Dielect. El. In. 2006, 13, 319-326.

(52) Viswanathan, G.; Chakrapani, N.; Yang, H.; Wei, B.; Chung, H.; Cho, K.; Ryu,
C. Y.; Ajayan, P. M. J. Am. Chem. Soc. 2003, 125, 9258-9259.

(53) Lewis, T. J. J. Phys. D. Appl. Phys. 1994, 38, 202-212.

(54) Nelson, J. K.; Fothergill, J. C. Nanotechnology 2004, 15, 586-595.

(55) Roy, M.; Nelson, J. K.; MacCrone, R. K.; Schadler, L. S. J. Mater. Sci. 2007, 42,
3789-3799.

(56) Cao, Y.; Irwin, P. C.; Younsi, K. IEEE T. Dielect. El. In. 2004, 11, 797-807.

90

Chapter 5. Size- and Ligand-Dependent Dielectric Properties of
pDCPD/BaTiO
3
Nanocomposites

5.1 Introduction
Recently, there has been significant interest in the development of capacitors that can
meet the needs of pulsed power energy storage systems requiring high breakdown
strength, low dielectric loss, and high thermal stability. By using a nanocomposite
approach, it has become possible to take advantage of the solution processability, high
dielectric strength, and low loss of polymers combined with the high permittivity of
inorganic nanocrystals to achieve these goals. The maximum energy density for a
capacitor is given by U
v
= ½ ε
r
ε
0
E
BD
2
J cm
-3
, where ɛ
r
is the relative permittivity of the
nanocomposite, ε
0
is the permittivity of free space, and E
BD
is the breakdown strength of
the nanocomposite. BaTiO
3
has gained significant interest as a filler material in
nanocomposites, having bulk permittivities as high as ε
r
= 10,000.
1,2
While the addition
of nanocrystalline BaTiO
3
into polymer matrices has been shown to systematically
increase composite permittivity, its presence often results in a cataclysmic reduction of
the nanocomposite breakdown strength, lessening any benefit achieved from increased
permittivity in terms of energy density.
3–6

In situ polymerization has been recently explored as a viable option in obtaining high
energy density BaTiO
3
-based nanocomposites.
7
By homogenizing the nanocrystals with
the monomer(s) prior to polymerization, it is possible to enhance the interaction between
the nanocrystal and polymer, reducing agglomeration and improving dielectric

91

performance.
8,9
A wide range of innovative techniques have been developed employing
both tailored nanocrystal and polymer chemistries to optimize devices, but no consistent
trend has emerged.
9–18
Despite the added level of control with in situ polymerization, the
nature of the polymer (i.e. purity, morphology, molecular weight) or its precise
interaction with BaTiO
3
is difficult to predict or measure. We have recently shown that
the presence of unmodified Ba
0.7
Sr
0.3
TiO
3
nanocrystals reduced polymer chain mobility
and increased the glass transition temperature, preventing the complete imidization of a
polyimide system polymerized in situ.
19
It is therefore of interest to develop a system
which can reduce complicating factors and overcome the aforementioned limitations
associated with in situ nanocomposite fabrication and gain a better understanding of the
direct effects of BaTiO
3
on nanocomposite performance.
4,6,20–22

Polydicyclopentadiene (pDCPD) is a crosslinked thermoset polymer that can be prepared
via ring opening metathesis polymerization (ROMP) of dicyclopentadiene (DCPD) with
the use of ruthenium-based Grubbs catalysts.
23–25
The polymerization reaction is done in
neat monomer (solvent-free), is highly exothermic, and operates with very small catalyst
loadings (< 1 wt%). Additionally, pDCPD has found widespread application as a result
of its high thermal stability, chemical resistance, low water uptake, and excellent
mechanical strength.
24,26
In comparison to the well-known thermoplastics polypropylene
and polystyrene, pDCPD has similar dielectric properties but is thermally stable up to 500
°C, extending its use to a variety of high temperature applications that are otherwise
unobtainable in polypropylene and polystyrene.
27
While several studies have explored
the physical and mechanical aspects of pDCPD nanocomposites,
28–30
there has only been

92

one report on the dielectric properties of a pDCPD nanocomposite.
31
Yin and coworkers
recently published the first known study on the dielectric properties of pDCPD-based
nanocomposites and discovered that pure pDCPD thin films possess a breakdown
strength (E
BD
) as high as 750 V μm
-1
, a modest dielectric constant (ε
r
= 2.5), and very low
dielectric loss (tan δ = 0.01). It was found that by adding 10 wt% of 20 nm fumed SiO
2

to the polymer, the high voltage coronal resistance improved and the permittivity of the
composite increased; however, the breakdown strength of resulting nanocomposites was
not reported.
31

Nanocomposites of well-dispersed BaTiO
3
nanocrystals in pDCPD have the potential to
harness the dielectric properties of BaTiO
3
while retaining some of the beneficial
properties of pDCPD; however, to date there has been no investigation into the use of
BaTiO
3
nanocrystals as a filler in pDCPD nanocomposites. In this study, we present the
first example of pDCPD/BaTiO
3
nanocomposites via an in situ polymerization route.
Commercially available 50 and 100 nm BaTiO
3
nanocrystals (BT50 and BT100,
respectively) were used to study the size-dependent dielectric properties of pDCPD-based
nanocomposites, while surface modification by the addition of organic ligands to the
nanocrystalline fillers was used to affect the nanocrystal-polymer interface and improve
device performance. The effect of filler size and surface modification in the
pDCPD/BaTiO
3
nanocomposite system was evaluated at 5 vol% (25 wt%) nanocrystal
loading, and results indicated a distinct ligand- and size-dependent effect on the measured
breakdown strength. To the best of our knowledge, this system has the highest reported
breakdown strength for any BaTiO
3
-based nanocomposite.

93

5.2 Experimental Section
All manipulations were performed under a nitrogen atmosphere using dry, air-free
solvents throughout. 50 nm BaTiO
3
nanocrystals (BT50) ( ≥ 99.9% trace metals basis, 50
nm by SEM; U.S. Research Nanomaterials, Inc.), 100 nm BaTiO
3
nanocrystals (BT100)
(99.0% trace metals basis, 100 nm by BET; Sigma Aldrich), 10-undeceneoic acid (99%;
Alfa Aesar), dicyclopentadiene (DCPD) (95%; Alfa Aesar), [1,3-bis(2,4,6-
trimethylphenyl)-2-imidazolidinylidene]dichloro(phenylmethylene)
(tricyclohexylphosphine)ruthenium (Grubbs catalyst, second generation; Materia, Inc.),
potassium permanganate (98%; Alfa Aesar), oxalic acid dihydrate (98%; Alfa Aesar),
and concentrated sulfuric acid (EMD Chemicals, Inc.) were all used as received.
5.2.1 Surface Functionalization of BaTiO
3
with 10-Undeceneoic Acid
In order to surface modify the nanocrystals, 1.58 g (8.60 mmol) of 10-undeceneoic acid
was dissolved in 50 mL of dry toluene and added directly to a 100 mL round bottom flask
containing 1.0 g (4.3 mmol) of BT50 or BT100 nanocrystals. The flask was placed under
flowing nitrogen and immersed into an ultrasonic bath. The solution was sonicated for 1
h at 30 °C before the nanocrystals were isolated via centrifugation (6000 rpm for 15 min).
The resulting solid was washed in toluene (3  30 mL) before being dried in vacuo (20
°C, 0.05 mmHg) for 24 h, and stored in a dry, air-free environment.

94

5.2.2 Material Characterization
Transmission electron microscopy (TEM) analysis was performed on a JEOL JEM-2100
microscope at an operating voltage of 200 kV, which was equipped with a Gatan Orius
CCD camera. Samples for TEM analysis were prepared by drop casting the nanocrystals
from toluene onto ultrathin carbon film supported on 400 mesh copper grids (Ted Pella,
Inc.). Scanning Electron Microscopy (SEM) analysis was performed on a JEOL JSM-
6610LV instrument in high-vacuum mode using an accelerating voltage of 10 kV. Cross-
section samples for SEM were prepared by freeze-fracturing a film, pressing it between
two glass plates and mounting on top of an aluminum stub. To prevent film charging, a
thin layer of carbon was deposited on all films. Multi-point BET measurements were
performed on a Nova 2200e surface area and pore size analyzer (Quantachrome
Instruments, Inc.). Fourier transform infrared (FT-IR) spectra were collected on a Bruker
Vertex 80v at a scanning interval of 0.5 cm
-1
with a resolution of 2 cm
-1
under flowing
nitrogen. Powder X-Ray diffraction (XRD) was performed on a Rigaku Ultima IV
diffractometer using Cu Kα (λ = 1.54 Å) radiation. For XRD samples of nanocrystalline
powder and pDCPD/BaTiO
3
films, the powder or film was placed directly onto a zero
diffraction single crystalline silicon substrate (MTI Corporation, Inc.).
Thermogravimetric analysis (TGA) was performed on a TA instruments Q50
thermogravimetric analyzer under flowing nitrogen at a heating rate of 10 °C min
-1
. All
samples for TGA analysis were dried within the instrument for 30 min under flowing
nitrogen at 100 °C prior to analysis. Dielectric constant and loss tangent were measured
on an Agilent 4294A impedance analyzer with a frequency sweep of 1 kHz – 2 MHz

95

using a parallel-plate plate configuration (plate area = 17 mm
2
). Dc negative polarity
breakdown measurements were performed on a Bertan 225, 20 kV dc power supply and
were monitored by an oscilloscope (Tektronix TDS 2004C; Tektronix, Inc.).
5.2.3 Device Preparation
The pDCPD/BaTiO
3
nanocomposites were prepared via an in situ polymerization route.
In a typical experiment, BaTiO
3
nanocrystals (variable amounts depending upon surface
ligand coverage) were added to 1.0 g DCPD. The DCPD was heated to 35-40°C to melt
the monomer, and then the mixture was immersed into an ultrasonic bath for 1 h with
intermittent mechanical mixing throughout. The water temperature in the ultrasonic bath
was kept at 35 °C to ensure the DCPD monomer remained melted and at low viscosity.
To initiate polymerization, 2 mg of second generation Grubbs catalyst (0.2 wt%) was
dissolved in 0.1 mL dry dichloromethane immediately prior to being added directly to the
monomer/nanocrystal mixture. The solution was continually mixed until a slight increase
in viscosity was evident (ca. 1 min), and then 0.4 mL of the reaction mixture was
removed and immediately cast between two 6”  6” cleaned glass plates and pressed with
50 N of force for 1 h at 20 °C. The pressed film was then placed into an oven at 100 °C
for 24 h under flowing nitrogen to ensure complete polymerization. The resulting film
thickness was controlled by placing 25.4 μm thick polymer shims (Practi-Shim; Accutrex
Products, Inc.) between the glass plates prior to casting. The film was slowly cooled to
room temperature under flowing nitrogen for 2 h before removing from the oven and
lifting the film. Upon reaching room temperature, the free standing film was lifted from

96

the glass plates by immersing into deionized water. Upon lifting, the film was blotted dry
and cut into squares 1.27  1.27 cm
2
in size and stored under nitrogen. Film thickness
ranged between 20-30 μm as measured by a custom metrology tool with ±1 μm accuracy.
Glass plates were cleaned by immersing in a very dilute solution of KMnO
4
/H
2
SO
4
(1:1
approximate molar ratio) for a minimum of 48 h. The plates were removed from the
solution, rinsed with deionized water, and then soaked in a dilute aqueous solution of
oxalic acid for 30 min. The plates were removed from the oxalic acid bath, rinsed with
deionized water, dried under flowing nitrogen, and immediately used.
5.2.4 Device Characterization
In order to measure the capacitance and loss tangent of the as-prepared free standing
films, a custom built dielectric testing station comprised of flat circular aluminum
electrodes was used. The film was placed directly between the two electrodes and
pressed, forming a parallel plate geometry (17 mm
2
plate area). Capacitance and loss
tangents were measured on an Agilent 4294A impedance analyzer with a frequency
sweep of 1 kHz – 2 MHz. Dielectric breakdown measurements were performed using a
gold plated hemispherical rod, 1 mm in diameter, which made direct contact with the
pDCPD/BaTiO
3
film. Upon contact, the film was immersed into a mineral oil bath then
subjected to a dc negative polarity bias at a linear ramp rate of 2500 V s
-1
(Bertan 225, 20
kV dcpower supply), connected to a high-voltage probe (Tektronix P6015A; Tektronix,
Inc.), and monitored by an oscilloscope (Tektronix TDS 2004C; Tektronix, Inc).
Breakdown was indicated by a spontaneous increase in current. One breakdown test was

97

performed per 1.6 cm
2
square, testing at least 15-20 squares for each nanocomposite.
The points for breakdown tests were chosen at random, and the thickness was measured
after each breakdown event within close proximity to the failure site.
5.3 Results and Discussion
5.3.1 Surface Modification and Characterization of BaTiO
3
Nanocrystals
The BT50 or BT100 nanocrystals were sonicated with 10-undeceneoic acid in dry toluene
to give the surface-modified nanocrystals, hereafter referred to as mBT50 and mBT100,
respectively. Qualitatively, the nanocrystals become more dispersible in nonpolar
organic solvents (e.g., toluene) after surface modification. Both mBT50 and mBT100
were stable throughout extensive rinsing with toluene, suggesting a strong binding
interaction between the 10-undeceneoic acid and the nanocrystal surface. After washing,
there remained ca. 1 carboxylate group per 6 nm
2
of BaTiO
3
surface area, assuming an

Figure 5.1. Powder XRD pattern of (a) BT50 powders and pDCPD/BT50 nanocomposite films
and (b) BT100 powders and pDCPD/BT100 nanocomposite films.

98

area of ~0.21 nm
2
for each carboxylate group, verified by thermogravimetric analysis
(TGA) and surface area measurements. Powder X-ray diffraction (XRD) of the modified
nanocrystals reveal that they are unaffected by the surface treatment and remain in the
cubic perovskite phase without the creation of any impurities phases such as BaCO
3
, with
a measured lattice constant of a = 4.04 ± 0.05 Å (JCPDS no. 75-0215) (Figure 5.1a,b).
Functionalization of mBT50 and mBT100 was confirmed by Fourier transform infrared
spectroscopy (FT-IR), as indicated by the presence of strong aliphatic v(C—H)
symmetric and asymmetric stretching bands between 2960 and 2860 cm
-1
and a weaker
primary alkene v(C—H) band at 3080 cm
-1
that were not present in either BT50 or
BT100 (Figure 5.2). Compared to traditional modification techniques which require long
reaction times and high temperatures to effectively modify the surface of nanocrystals,
sonication allowed for functionalization to occur at lower temperatures over a shorter
period of time.
32–35
In addition, sonication effectively broke apart larger agglomerated

Figure 5.2. FT-IR spectra of BT50, BT100, mBT50, and mBT100 nanocrystals.

99

particles and exposed more surface area for increased reactivity and improved dispersion.
5.3.2 pDCPD/BaTiO
3
Nanocomposite Fabrication and Characterization
Devices for dielectric testing were prepared via a solvent-free in situ polymerization route
whereby 5 vol% (25 wt%) of BT50, BT100, mBT50, or mBT100 powders were added
directly to the DCPD monomer. The polymerization was started by addition of 0.2 wt%
second generation Grubbs catalyst pre-dissolved in dichloromethane (DCM) to produce
free-standing thermoset films. The resulting sheet-like films were between 20 – 30 μm
thick. The pure pDCPD film was extremely flexible, optically clear, and displayed a
faint yellow hue as a result of the ruthenium-based catalyst. Upon addition of 5 vol%
BT50, BT100, mBT50, or mBT100, the nanocomposite became opaque yet remained

Figure 5.3. Digital photograph of pDCPD/BT100 demonstrating film flexibility.

100

extremely flexible and could be handled without cracking or breaking (Figure 5.3). The
addition of nanocrystals had no noticeable effect on the nanocomposite processability as
a result of the very low viscosity of the DCDP monomer, which ensured effective mixing.
XRD analysis of the nanocomposite films revealed that the BaTiO
3
nanocrystal fillers
remained phase pure after processing, with no indication of BaCO
3
formation (Figure
5.1). Scanning electron microscopy (SEM) images of the pDCPD/BT50 nanocomposites
revealed relatively uniform film morphology with moderate agglomeration and particle
settling, whereas the pDCPD/BT100 system had larger agglomerates and more
significant settling (Figure 5.4a,b). SEM images of pDCPD/mBT50 and
pDCPD/mBT100 revealed improved dispersion over the nanocomposites containing

Figure 5.4. SEM micrographs of (a) pDCPD/BT50, (b) pDCPD/BT100, (c) pDCPD/mBT50, and
(d) pDCPD/mBT100.

101

unmodified BaTiO
3
, displaying smaller agglomerates and more consistent dispersion over
the film area (Figure 5.4c,d). The presence of 10-undecenoic acid in mBT50 and
mBT100 should screen the highly polar BaTiO
3
nanocrystal surface and created a more
favorable interaction at the nanoparticle-polymer interface, thereby improving dispersion
within the polymer matrix and positively affecting device performance.
The pDCPD/BT composites were thermally robust, as confirmed by TGA (Figure
5.5a,b). Addition of the BT50 or BT100 fillers had no effect on the thermal stability
relative to the neat pDCPD, which possessed a decomposition onset at 450 °C under
nitrogen (as determined by the onset of the first derivative of the mass loss as a function
of temperature). The pDCPD/mBT50 and pDCPD/mBT100 nanocomposites displayed

Figure 5.5. TGA curves in nitrogen for all (a) pDCPD/BT50 and (b) pDCPD/BT100
nanocomposites at 5 vol% loading, and all (c) BT50 and (d) BT100 nanocrystals.

102

slightly reduced thermal stability, with decomposition onsets occurring at 390 °C. Low
temperature mass loss was attributed to decomposition of the strongly bound ligands at
the nanocrystal surface, and is consistent with the decomposition onset of thermograms
obtained for mBT50 and mBT100 (Figure 5.5 c,d).
5.3.3 Dielectric Characterization
The dielectric constant and loss tangent of the nanocomposites were measured at a range
of frequencies from 1 kHz to 2 MHz. The dielectric constant of the pDCPD/BaTiO
3

nanocomposites was higher than neat pDCPD at all frequencies tested. For example, the
relative permittivity of pDCPD was ε
r
= 1.7 at 1.0 MHz and increased to ε
r
= 2.4 after the
addition of 5 vol% BT50 (Figure 5.6a). A dielectric constant between 2.3 and 2.4 was
measured for nanocomposites with both crystallite sizes, being only marginally lower in
the case of pDCPD/mBT50. In all of the nanocomposites tested, the dielectric loss
remained below tan δ = 0.008 from 100 kHz – 2 MHz , and was improved over the neat
pDCPD, which displayed dielectric loss below 0.013 over the same rang(Figure 5.6b).
Maintaining low loss in dielectric films is desirable as it reduces the likelihood for
leakage currents that facilitate premature breakdown at low electric fields.
36
The
dielectric loss is in close agreement with Yin and coworkers,
31
and is markedly improved
over PVDF systems that have shown dielectric losses as high as 0.17 with 5 vol% loading
of 100 nm BaTiO
3
(at 1 MHz).
6

103

To study the effect of nanocrystal size and ligands on the dielectric breakdown strength
of pDCPD nanocomposites, free-standing films were subjected to a negative dc bias until
breakdown was indicated by an instantaneous increase in current. Dielectric breakdown
strength was calculated using a two-parameter Weibull distribution function:
36
P = 1-
exp[-(E/E
BD
)
ß
], where P is the cumulative probability of breakdown, E is the
experimentally measured breakdown strength, E
BD
is the cumulative breakdown
probability at 63.2%, and β is a shape parameter for 15-20 randomly selected independent
measurements per composite. The Weibull distributions of breakdown strengths are
given in Figure 5.7. The breakdown strength of the pDCPD/BaTiO
3
nanocomposites
increases as the size of the nanocrystals decrease, reaching E
BD
= 405 V μm
-1
for
pDCPD/BT50. This breakdown strength is 31% greater than that for pDCPD/BT100,
which possessed E
BD
= 308 V μm
-1
. The increase in breakdown strength is consistent
with BT50 having a higher surface area (27 m
2
g
-1
for BT50 compared to 17 m
2
g
-1
for

Figure 5.6. (a) Relative permittivity and dielectric loss of pDCPD, pDCPD/BT50, and
pDCPD/mBT50 devices as a function of frequency. (b) Relative permittivity and dielectric loss
of pDCPD, pDCPD/BT100, and pDCPD/mBT100 devices as a function of frequency.
Measurements were taken at 25°C.

104

BT100) and a more homogeneous dispersion within the polymer matrix than BT100 at
the same loading (vide supra). The increased surface area for BT50 allows for increased
interfacial interaction between the nanocrystal and polymer.
37,38
This short-range
interfacial zone surrounding the nanoparticulate fillers can potentially account for a
significant volume fraction of the composite, dominating the dielectric response as the
size of the nanocrystal decreases.
39,40
The absolute size of the interfacial zone is largely
unknown, but has been shown to contribute to charge dissipation and suppression of
interfacial polarization in smaller nanoparticulate fillers.
40,41
Surface modification with
10-undeceneoic acid further improved the high voltage endurance with both mBT100 and
mBT50 nanocrystals by up to 20%, reaching a maximum E
BD
= 468 V μm
-1
for
pDCPD/mBT50. Through surface modification, it was possible to tune the interfacial
interaction at the nanocrystal surface and further enhance the composite dielectric
properties. The presence of a non-polar organic corona surrounding the nanocrystal
surface enhanced the size effects by creating an insulating barrier that more closely
matched the polymer polarity and mitigated localized charge accumulation and reduced
charge carrier mobility, facilitating an increase in the breakdown strength of both
systems.
42,43

As a result of the improved breakdown strength and increased dielectric constant (ε
r
=
2.4) in pDCPD/mBT50, the calculated energy density (2.4 J cm
-3
) rose by 5% over the
pure polymer. While the breakdown strength of pDCPD (E
BD
= 541 V μm
-1
) was greater
than pDCPD/BT50, the larger relative permittivity of the 5 vol% nanocomposite was
sufficient to provide a higher calculated energy density.

105

Traditionally, BaTiO
3
-based nanocomposites display a significant decrease in breakdown
strength at very low loadings and only achieve improved energy densities with high
concentrations of filler.
44
Dang, Feng and coworkers prepared polyimide/BaTiO
3

nanocomposites in situ and experienced a 40 % decrease in breakdown strengths at
loadings as low as 5 vol%, with improved energy densities only achieved at loadings
exceeding 30 vol%.
3,12
At such high loadings, effective medium models predict
significant improvement in the dielectric constant, which overcomes losses associated
with reduced breakdown strength, thereby creating systems in which the filler material
dominates the overall dielectric response.
45,46
In this system, surface modification has
effectively been shown to mitigate substantial drops in breakdown strength at low
loadings through increased interfacial interaction and improved dispersion.

Figure 5.7. Two-parameter Weibull plots of pDCPD/BaTiO
3
nanocomposites where the dashed
line represents Weibull breakdown at 63.2 % probability of failure. Measurements were taken
at 25 °C.

106

5.4 Conclusions
Nanocrystals of BaTiO
3
were successfully blended into pDCPD via an in-situ
polymerization route to understand the effects filler size and surface ligands on the
dielectric properties of the nanocomposite. At 5 vol% loading BaTiO
3
, the relative
permittivity of the resulting pDCPD/BaTiO
3
composites increased from ε
r
= 1.7 for the
pure polymer to a maximum of ε
r
= 2.4 for the composite, while tan δ < 0.008 was
maintained for all compositions up to 2 MHz. The breakdown strength for the
nanocomposites with no surface modification demonstrated a pronounced size
dependence, with the smaller BT50 crystallites having a breakdown strength of 405 V
μm
-1
, a 31% increase over nanocomposites containing BT100 at the same loading. By
modifying the surface of the nanocrystals, it was possible to further improve the
BaTiO
3
—pDCPD interfacial effects, increasing the breakdown voltage to 468 V μm
-1
for
the pDCPD/mBT50 nanocomposite. As a result of surface modification, the measured
energy density was increased over the pure polymer, highlighting the potential of this
system for the future low loss capacitors requiring high thermal stability.
5.5 References
(1) Lines, M. E. Principles and applications of ferroelectrics and related materials;
Oxford classic texts in the physical sciences; Clarendon Press ; Oxford University
Press: Oxford : New York, 2001.

(2) Newnham, R. E.; Cross, L. E. MRS Bull. 2011, 30, 845–848.

(3) Dang, Z.-M.; Lin, Y.-Q.; Xu, H.-P.; Shi, C.-Y.; Li, S.-T.; Bai, J. Adv. Funct.
Mater. 2008, 18, 1509–1517.

107

(4) Kim, P.; Jones, S. C.; Hotchkiss, P. J.; Haddock, J. N.; Kippelen, B.; Marder, S.
R.; Perry, J. W. Adv. Mater. 2007, 19, 1001–1005.

(5) Barber, P.; Balasubramanian, S.; Anguchamy, Y.; Gong, S.; Wibowo, A.; Gao,
H.; Ploehn, H. J.; Zur Loye, H.-C. Materials 2009, 2, 1697–1733.

(6) Kim, P.; Doss, N. M.; Tillotson, J. P.; Hotchkiss, P. J.; Pan, M.-J.; Marder, S. R.;
Li, J.; Calame, J. P.; Perry, J. W. ACS Nano 2009, 3, 2581–2592.

(7) Chen, B.; Evans, J. R. G.; Greenwell, H. C.; Boulet, P.; Coveney, P. V.; Bowden,
A. A.; Whiting, A. Chem. Soc. Rev. 2008, 37, 568.

(8) Nanocomposites: preparation, properties, and performance; Nova Science
Publishers: New York, 2008.

(9) Friedrich, K. Polymer composites: from nano-to-macro-scale; Springer: New
York, 2005.

(10) Patsidis, A. C.; Kalaitzidou, K.; Psarras, G. C. Mater. Chem. Phys. 2012, 135,
798–805.

(11) Tomer, V.; Polizos, G.; Manias, E.; Randall, C. A. J. Appl. Phys. 2010, 108,
074116.

(12) Yu Feng; Jinghua Yin; Minghua Chen; Xiaoxu Liu; Guang Li IEEE, 2011; pp.
226–229.

(13) Choudhury, A. Mater. Chem. Phys. 2010, 121, 280–285.

(14) Xie, S.-H.; Zhu, B.-K.; Wei, X.-Z.; Xu, Z.-K.; Xu, Y.-Y. Compos. Part A.-Appl.
S. 2005, 36, 1152–1157.

(15) Siddabattuni, S.; Schuman, T. P.; Dogan, F. Mater. Sci. Eng. B. 2011, 176, 1422–
1429.

(16) Xie, L.; Huang, X.; Wu, C.; Jiang, P. J. Mater. Chem. 2011, 21, 5897.

(17) Yang, K.; Huang, X.; Xie, L.; Wu, C.; Jiang, P.; Tanaka, T. Macromol. Rapid
Comm. 2012, 33, 1921–1926.

(18) Chon, J.; Ye, S.; Cha, K. J.; Lee, S. C.; Koo, Y. S.; Jung, J. H.; Kwon, Y. K.
Chem. Mater. 2010, 22, 5445–5452.

108

(19) Beier, C. W.; Sanders, J. M.; Brutchey, R. L. J. Phys. Chem. C 2013,
130221152520004.

(20) Neouze, M.-A.; Schubert, U. Montash. Chem. 2008, 139, 183–195.

(21) Almadhoun, M. N.; Bhansali, U. S.; Alshareef, H. N. J. Mater. Chem. 2012, 22,
11196.

(22) Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Langmuir 2010, 26, 5067–5071.

(23) Slugovc, C. Macromol. Rapid Comm. 2004, 25, 1283–1297.

(24) Handbook of metathesis; Wiley-VCH: Weinheim, Germany, 2003.

(25) Kissin, Y. V. In Kirk-Othmer Encyclopedia of Chemical Technology; John Wiley
& Sons, Inc., Ed.; John Wiley & Sons, Inc.: Hoboken, NJ, USA, 2005.

(26) Grubbs, R. H. Tetrahedron 2004, 60, 7117–7140.

(27) Rabuffi, M.; Picci, G. IEEE T. Plasma Sci. 2002, 30, 1939–1942.

(28) Simons, R.; Guntari, S. N.; Goh, T. K.; Qiao, G. G.; Bateman, S. A. J. Polym. Sci.
Part A. 2012, 50, 89–97.

(29) Jeong, W.; Kessler, M. R. Chem. Mater. 2008, 20, 7060–7068.

(30) Yoonessi, M.; Toghiani, H.; Kingery, W. L.; Pittman, C. U. Macromolecules
2004, 37, 2511–2518.

(31) Yin, W.; Kniajanski, S.; Amm, B. IEEE, 2010; pp. 1–5.

(32) Xu, C.; Xu, K.; Gu, H.; Zheng, R.; Liu, H.; Zhang, X.; Guo, Z.; Xu, B. Journal of
the American Chemical Society 2004, 126, 9938–9939.

(33) Shultz, M. D.; Reveles, J. U.; Khanna, S. N.; Carpenter, E. E. Journal of the
American Chemical Society 2007, 129, 2482–2487.

(34) Huang, W.; Lin, Y.; Taylor, S.; Gaillard, J.; Rao, A. M.; Sun, Y.-P. Nano Letters
2002, 2, 231–234.

(35) Koshio, A.; Yudasaka, M.; Zhang, M.; Iijima, S. Nano Letters 2001, 1, 361–363.

(36) Dissado, L. A. Electrical degradation and breakdown in polymers; IEE materials
and devices series; P. Peregrinus: London, 1992.

109

(37) Lewis, T. J. J. Phys. D: Appl. Phys. 2005, 38, 202–212.

(38) Viswanathan, G.; Chakrapani, N.; Yang, H.; Wei, B.; Chung, H.; Cho, K.; Ryu,
C. Y.; Ajayan, P. M. J. Am. Chem. Soc. 2003, 125, 9258–9259.

(39) Nelson, J. K.; Fothergill, J. C. Nanotechnology 2004, 15, 586–595.

(40) Roy, M.; Nelson, J. K.; MacCrone, R. K.; Schadler, L. S.; Reed, C. W.; Keefe, R.;
Zenger, W. IEEE T. Dielect. El. In. 2005, 12, 629–643.

(41) Roy, M.; Nelson, J. K.; MacCrone, R. K.; Schadler, L. S. J. Mater. Sci. 2007, 42,
3789–3799.

(42) Jung, H. M.; Kang, J.-H.; Yang, S. Y.; Won, J. C.; Kim, Y. S. Chem. Mater. 2010,
22, 450–456.

(43) Ma, D.; Siegel, R. W.; Hong, J.-I.; Schadler, L. S.; Mårtensson, E.; Önneby, C. J.
Mater. Res. 2011, 19, 857–863.

(44) Khalil, M. S. IEEE T. Dielect. El. In. 2000, 7, 261–268.

(45) An, A. L.; Boggs, S. A.; Calame, J. IEEE Electr. Insul. M. 2008, 24, 5–10.

(46) An, A. L.; Boggs, S. A.; Calame, J. IEEE Int. Symp. El. In. 2006, 466–469.

110
Bibliography
Alam, M. A.; Azarian, M. H.; Pecht, M. G. J. Mater. Sci.-Mater. El. 2012, 23, 1504–
1510.
Allara, D. L.; Nuzzo, R. G. Langmuir 1985, 1, 52–66.
Almadhoun, M. N.; Bhansali, U. S.; J. Mater. Chem. 2012, 22, 11196.
Althues, H.; Henle, J.; Kaskel, S. Chem. Soc. Rev. 2007, 36, 1454.
An, A. L.; Boggs, S. A.; Calame, J. IEEE Int. Symp. El. In. 2006, 466–469.
An, A. L.; Boggs, S. A.; Calame, J. IEEE Electr. Insul. M. 2008, 24, 5–10.
Arya, P. R.; Jha, P.; Ganguli, A. K. J. Mater. Chem. 2003, 13, 415–423.
Ash, B. J.; Siegel, R. W.; Schadler, L. S. J. Polym. Sci. Pol. Phys. 2004, 42, 4371–4383.
Bansal, V.; Poddar, P.; Ahmad, A.; Sastry, M. J. Am. Chem. Soc. 2006, 128, 11958–
11963.
Barber, P.; Balasubramanian, S.; Anguchamy, Y.; Gong, S.; Wibowo, A.; Gao, H.;
Ploehn, H. J.; Loye, H.-C. Materials 2009, 2, 1697–1733.
Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Small 2008, 4, 2102–2106.
Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. Langmuir 2010, 26, 5067–5071.
Beier, C. W.; Cuevas, M. A.; Brutchey, R. L. J. Mater. Chem. 2010, 20, 5074-5079.
Beier, C. W.; Sanders, J. M.; Brutchey, R. L. J. Phys. Chem. C 2013, 10.1021/jp312519r.
Bhalla, A. S.; Guo, R.; Roy, R. Mat. Res. Innovat. 2000, 4, 3–26.
Brandstetter, S. S.; Drummy, L. F.; Horwath, J. C.; Schweickart, D. L.; Vaia, R. A. In;
IEEE, 2008; pp. 287–290.
Brewer, S. H.; Brown, D. A.; Franzen, S. Langmuir 2002, 18, 6857–6865.
Brutchey, R. L.; Cheng, G.; Gu, Q.; Morse, D. E. Adv. Mater. 2008, 20, 1029–1033.
Brutchey, R. L.; Morse, D. E. Angew. Chem. Int. Ed. 2006, 45, 6564–6566.
Bunting, E. N.; Shelton, G. R.; Creamer, A. S. J. Am. Ceram. Soc. 1947, 30, 114–125.

111
Buscaglia, V.; Buscaglia, M. T.; Viviani, M.; Mitoseriu, L.; Nanni, P.; Trefiletti, V.;
Piaggio, P.; Gregora, I.; Ostapchuk, T.; Pokorný, J. J. Eur. Ceram. Soc. 2006, 26, 2889–
2898.
Calebrese, C.; Hui, L.; Schadler, L.; Nelson, J. IEEE T. Dielect. El. In. 2011, 18, 938–
945.
Caulton, K. G.; Hubert-Pfalzgraf, L. G. Chem. Rev. 1990, 90, 969–995.
Chandler, C. D.; Roger, C.; Hampden-Smith, M. J. Chem. Rev. 1993, 93, 1205–1241.
Chao, M.; Kou, K. C.; Wang, Z. C.; Wu, G. L.; Zhang, D. N.; Zhang, J. Q. Adv. Mater.
Res. 2011, 199-200, 13–18.
Chen, B.; Evans, J. R. G.; Greenwell, H. C.; Boulet, P.; Coveney, P. V.; Bowden, A. A.;
Whiting, A. Chem. Soc. Rev. 2008, 37, 568.
Chen, D.; Zhu, H.; Liu, T. ACS Appl. Mater. Interfaces 2010, 2, 3702–3708.
Chen, W.; Zhu, Q. Mater. Lett. 2007, 61, 3378–3380.
Cho, S.-D.; Lee, S.-Y.; Hyun, J.-G.; Paik, K.-W. J. Mater. Sci.-Mater. El. 2005, 16, 77–
84.
Chon, J.; Ye, S.; Cha, K. J.; Lee, S. C.; Koo, Y. S.; Jung, J. H.; Kwon, Y. K. Chem.
Mater. 2010, 22, 5445–5452.
Choudhury, A. Mater. Chem. Phys. 2010, 121, 280–285.
Chu, L. W.; Prakash, K. N.; Tsai, M.-T.; Lin, I.-N. J. Eur. Ceram. Soc. 2008, 28, 1205–
1212.
Coppard, R. W.; Bowman, J.; Dissado, L. A.; Rowland, S. M.; Rakowski, R. T. J. Phys.
D: Appl. Phys. 1990, 23, 1554–1561.
Dang, Z.-M.; Lin, Y.-Q.; Xu, H.-P.; Shi, C.-Y.; Li, S.-T.; Bai, J. Adv. Funct. Mater. 2008,
18, 1509–1517.
Dang, Z.-M.; Wang, H.-Y.; Peng, B.; Nan, C.-W. J. Electroceram. 2007, 21, 381–384.
, A. F.; Imhof, A. Chem. Mater. 2009, 21, 3002–3007.
Dissado, L. A. Electrical Degradation and Breakdown in Polymers; IEE materials and
devices series; P. Peregrinus: London, 1992.

112
Drobny, J. G.; Rapra Technology Limited Fluoroplastics; Rapra Technology Ltd.:
Shawbury, Shrewsbury, Shropshire, U.K., 2006.
Fernández-García, M.; Martínez-Arias, A.; Hanson, J. C.; Rodriguez, J. A. Chem. Rev.
2004, 104, 4063–4104.
Fillery, S. P.; Koerner, H.; Drummy, L.; Dunkerley, E.; Durstock, M. F.; Schmidt, D. F.;
Vaia, R. A. ACS Appl. Mater. Interfaces 2012, 4, 1388–1396.
Fragiadakis, D.; Logakis, E.; Pissis, P.; Kramarenko, V. Y.; Shantalii, T. A.; Karpova, I.
L.; Dragan, K. S.; Privalko, E. G.; Usenko, A. A.; Privalko, V. P. J. Phys. Conf. Ser.
2005, 10, 139–142.
Frey, M. H.; Payne, D. A. Phys. Rev. B. 1996, 54, 3158–3168.
Friedrich, K. Polymer Composites: From Nano-to-macro-scale; Springer: New York,
2005.
Gao, W.; Dickinson, L.; Grozinger, C.; Morin, F. G.; Reven, L. Langmuir 1996, 12,
6429–6435.
Gao, W.; Reven, L. Langmuir 1995, 11, 1860–1863.
Gao, Y.; Koumoto, K. Crystal Growth Des. 2005, 5, 1983–2017.
Gaskins, B. C.; Lannutti, J. J. J. Mater. Res. 2011, 11, 1953–1959.
Ge, J. J.; Xue, G.; Li, F.; McCreight, K. W.; Wang, S.-Y.; Harris, F. W.; Cheng, S. Z. D.;
Zhuang, X.; Hong, S.-C.; Shen, Y. R. Macromol. Rapid Comm. 1998, 19, 619–623.
Goldschmidt, V. M. Skrifter Norske Videnskaps-Akad. Oslo. I. Matemal.-Naturv. Klasse
1927, 8, 7–156.
Grubbs, R. H. Handbook of Metathesis; Wiley-VCH: Weinheim, Germany, 2003.
Grubbs, R. H. Tetrahedron 2004, 60, 7117–7140.
Hilton, A. D.; Ricketts, B. W. J. Phys. D: Appl. Phys. 1996, 29, 1321–1325.
Hornebecq, V.; Huber, C.; Maglione, M.; Antonietti, M.; Elissalde, C. Adv. Funct. Mater.
2004, 14, 899–904.
Hou, B.; Xu, Y.; Wu, D.; Sun, Y. Powder Technol. 2006, 170, 26–30.
Hou, R. Z.; Ferreira, P.; Vilarinho, P. M. Chem. Mater. 2009, 21, 3536–3541.

113
Hsiao, S.-H.; Chen, Y.-J. Eur. Polym. J. 2002, 38, 815–828.
Huang, L.; Chen, Z.; Wilson, J. D.; Banerjee, S.; Robinson, R. D.; Herman, I. P.;
La bow tz, R.; O’B n, S. J. Appl. Phys. 2006, 100, 034316.
Huang, W.; Lin, Y.; Taylor, S.; Gaillard, J.; Rao, A. M.; Sun, Y.-P. Nano Letters 2002, 2,
231–234.
Huber, C.; Elissalde, C.; Hornebecq, V.; Mornet, S.; Treguer-Delapierre, M.; Weill, F.;
Maglione, M. Ceram. Int. 2004, 30, 1241–1245.
Imai, T.; Sawa, F.; Nakano, T.; Ozaki, T.; Shimizu, T.; Kuge, S. -i.; Kozako, M.; Tanaka,
T. In; IEEE; pp. 171–174.
Ishida, H.; Wellinghoff, S. T.; Baer, E.; Koenig, J. L. Macromolecules 1980, 13, 826–
834.
Ishii, J.; Takata, A.; Oami, Y.; Yokota, R.; Vladimirov, L.; Hasegawa, M. Eur. Polym. J.
2010, 46, 681–693.
Israelachvili, J. N. Intermolecular and surface forces; Acad. Press: London [u.a.], 1997.
Jeong, W.; Kessler, M. R. Chem. Mater. 2008, 20, 7060–7068.
Jiang, B.; Peng, J. L.; Bursill, L. A. Ferroelectrics 1998, 207, 445–463.
Jiquan, H.; Maochun, H.; Feilong, J.; Yongge, C. Mater. Lett. 2008, 62, 2304–2306.
Jung, H. M.; Kang, J.-H.; Yang, S. Y.; Won, J. C.; Kim, Y. S. Chem. Mater. 2010, 22,
450–456.
Kao, K.-C. Dielectric Phenomena in Solids: With Emphasis on Physical Concepts of
Electronic Processes; Acad c P : A t da ; Bo ton, 2004.
Khalil, M. S. IEEE T. Dielect. El. In. 2000, 7, 261–268.
Kim, D.-W.; Lee, D.-H.; Kim, B.-K.; Je, H.-J.; Park, J.-G. Macromol. Rapid Comm.
2006, 27, 1821–1825.
Kim, P.; Jones, S. C.; Hotchkiss, P. J.; Haddock, J. N.; Kippelen, B.; Marder, S. R.;
Perry, J. W. Adv. Mater. 2007, 19, 1001–1005.
Kim, P.; Doss, N. M.; Tillotson, J. P.; Hotchkiss, P. J.; Pan, M.-J.; Marder, S. R.; Li, J.;
Calame, J. P.; Perry, J. W. ACS Nano 2009, 3, 2581–2592.

114
Kissin, Y. V. Polymers of Higher Olefins. In Kirk-Othmer Encyclopedia of Chemical
Technology; John Wiley & Sons, Inc., Ed.; John Wiley & Sons, Inc.: Hoboken, NJ, USA,
2005.
Kobayashi, Y.; Tanase, T.; Tabata, T.; Miwa, T.; Konno, M. J. Eur. Ceram. Soc. 2008,
28, 117–122.
Komarneni, S.; Parker, J. C.; Wollenberger, H. Nanophase and Nanocomposite Materials
II: Symposium Held December 2-5, 1996, Boston, Massachusetts, U.S.A; Materials
Research Society symposium proceedings; Materials Research Society: Pittsburgh, Pa,
1997.
Konieczny, J. M.; Wunder, S. L. Macromolecules 1996, 29, 7613–7615.
Koshio, A.; Yudasaka, M.; Zhang, M.; Iijima, S. Nano Letters 2001, 1, 361–363.
Lafond, V.; Gervais, C.; Maquet, J.; Prochnow, D.; Babonneau, F.; Mutin, P. H. Chem.
Mater. 2003, 15, 4098–4103.
Lewis, T. J. J. Phys. D: Appl. Phys. 2005, 38, 202–212.
Li, J.; Wu, Y. J.; Tanaka, H.; Yamamoto, T.; Kuwabara, M. J. Am. Chem. Soc. 2004, 87,
1578–1581.
Li, J.; Claude, J.; Norena-Franco, L. E.; Seok, S. I.; Wang, Q. Chem. Mater. 2008, 20,
6304–6306.
Lin, B.; Liu, H.; Zhang, S.; Yuan, C. J. Solid State Chem. 2004, 177, 3849–3852.
Lines, M. E. Principles and Applications of Ferroelectrics and Related Materials; Oxford
cla c t xt n th phy cal c nc ; Cla ndon P  ; Oxfo d Un v ty P : Oxfo d :
New York, 2001.
Ma, D.; Siegel, R. W.; Hong, J.-I.; Schadler, L. S.; Mårtensson, E.; Önneby, C. J. Mater.
Res. 2011, 19, 857–863.
Ma, D.; Hugener, T. A.; Siegel, R. W.; Christerson, A.; Mårtensson, E.; Önneby, C.;
Schadler, L. S. Nanotechnology 2005, 16, 724–731.
Mancini, L. H.; Esposito, C. L. Nanocomposites: Preparation, Properties, and
Performance; Nova Science Publishers: New York, 2008.
Mao, Y. P.; Mao, S. Y.; Ye, Z.-G.; Xie, Z. X.; Zheng, L. S. J. Appl. Phys. 2010, 108,
014102.
Mao, Y.; Banerjee, S.; Wong, S. S. Chem. Commun. 2003, 408–409.

115
Mao, Y.; Banerjee, S.; Wong, S. S. J. Am. Chem. Soc. 2003, 125, 15718–15719.
McMorrow, B.; Chartoff, R.; Lucas, P.; Richardson, W.; Anderson, P. Compos. Interface
2006, 13, 801–817.
Mendes, S. F.; Costa, C. M.; Caparros, C.; Sencadas, V.; Lanceros-Méndez, S. J. Mater.
Sci. 2011, 47, 1378–1388.
Merz, W. Phys. Rev. 1949, 75, 687–687.
Mittal, K. L. Polyimides and other high temperature polymers: synthesis,
characterization and applications; VSP: Savannah, GA, 2007.
Momen, G.; Farzaneh, M. Adv. Mater. Sci. 2011, 27, 1–13.
Murugaraj, P.; Mainwaring, D.; Mora-Huertas, N. J. Appl. Phys. 2005, 98, 054304.
Mutin, P. H.; Guerrero, G.; Vioux, A. J. Mater. Chem. 2005, 15, 3761-3768.
Nelson, J. K.; Fothergill, J. C. Nanotechnology 2004, 15, 586–595.
Neouze, M.-A.; Schubert, U. Montash. Chem. 2008, 139, 183–195.
Newnham, R. E.; Cross, L. E. Ferroelectricity: MRS Bull. 2011, 30, 845–848.
Niederberger, M.; Bartl, M. H.; Stucky, G. D. J. Am. Chem. Soc. 2002, 124, 13642–
13643.
Niederberger, M.; Garnweitner, G.; Pinna, N.; Antonietti, M. J. Am. Chem. Soc. 2004,
126, 9120–9126.
Niederberger, M.; Pinna, N.; Polleux, J.; Antonietti, M. Angew. Chem. Int. Ed. 2004, 43,
2270–2273.
O’B n, S.; B u , L.; Mu ay, C. B. J. Am. Chem. Soc. 2001, 123, 12085–12086.
Packia Selvam, I.; Kumar, V. Mater. Lett. 2002, 56, 1089–1092.
Paniagua, S. A.; Hotchkiss, P. J.; Jones, S. C.; Marder, S. R.; Mudalige, A.; Marrikar, F.
S.; Pemberton, J. E.; Armstrong, N. R. J. Phys. Chem. C 2008, 112, 7809–7817.
Pant, H.; Patra, M.; Verma, A.; Vadera, S.; Kumar, N. Acta Mater. 2006, 54, 3163–3169.
Patsidis, A. C.; Kalaitzidou, K.; Psarras, G. C. Mater. Chem. Phys. 2012, 135, 798–805.
Pinna, N.; Niederberger, M. Angew. Chem. Int. Ed. 2008, 47, 5292–5304.

116
Pramanik, N. C.; Anisha, N.; Abraham, P. A.; Panicker, N. R. J. Alloy. Compd. 2009,
476, 524–528.
Rabuffetti, F. A.; Brutchey, R. L. Chem. Commun. 2012, 48, 1437.
Rabuffetti, F. A.; Brutchey, R. L. J. Am. Chem. Soc. 2012, 134, 9475–9487.
Rabuffetti, F. A.; Lee, J. S.; Brutchey, R. L. Chem. Mater. 2012, 24, 3114–3116.
Rabuffi, M.; Picci, G. IEEE T. Plasma Sci. 2002, 30, 1939–1942.
Roy, M.; Nelson, J. K.; MacCrone, R. K.; Schadler, L. S.; Reed, C. W.; Keefe, R.;
Zenger, W. IEEE T. Dielect. El. In. 2005, 12, 629–643.
Roy, M.; Nelson, J. K.; MacCrone, R. K.; Schadler, L. S. J. Mater. Sci. 2007, 42, 3789–
3799.
Shultz, M. D.; Reveles, J. U.; Khanna, S. N.; Carpenter, E. E. Journal of the American
Chemical Society 2007, 129, 2482–2487.
Siddabattuni, S.; Schuman, T. P.; Dogan, F. Mater. Sci. Eng. B. 2011, 176, 1422–1429.
Simons, R.; Guntari, S. N.; Goh, T. K.; Qiao, G. G.; Bateman, S. A. J. Polym. Sci. Part A.
2012, 50, 89–97.
Sing, K. S. W. Pure Appl. Chem. 1985, 57, 603–619.
Slugovc, C. Macromol. Rapid Comm. 2004, 25, 1283–1297.
Smith, M. B.; Page, K.; Siegrist, T.; Redmond, P. L.; Walter, E. C.; Seshadri, R.; Brus, L.
E.; Steigerwald, M. L. J. Am. Chem. Soc. 2008, 130, 6955–6963.
Song, Y.; Shen, Y.; Liu, H.; Lin, Y.; Li, M.; Nan, C.-W. J. Mater. Chem. 2012, 22,
16491.
Spaldin, N. A. Science 2004, 304, 1606–1607.
Spanier, J. E.; Kolpak, A. M.; Urban, J. J.; Grinberg, I.; Ouyang, L.; Yun, W. S.; Rappe,
A. M.; Park, H. Nano Lett. 2006, 6, 735–739.
Sroog, C. E. J. Polym. Sci. Macromol. Rev. 1976, 11, 161–208.
Sroog, C. E.; Endrey, A. L.; Abramo, S. V.; Berr, C. E.; Edwards, W. M.; Olivier, K. L.
J. Polym. Sci. Part A. 1965, 3, 1373–1390.

117
Stuart, B. Infrared Spectroscopy: Fundamentals and Applications; Analytical techniques
n th c nc ; J. W l y: Ch ch t , W t Su x, England ; Hobok n, NJ, 2004.
Su, K.; Nuraje, N.; Yang, N.-L. Langmuir 2007, 23, 11369–11372.
Takeuchi, T.; Tabuchi, M.; Ado, K.; Honjo, K.; Nakamura, O.; Kageyama, H.; Suyama,
Y.; Ohtori, N.; Nagasawa, M. J. Mater. Sci. 1997, 32, 4053–4060.
Tamai, S.; Yamaguchi, A.; Ohta, M. Polymer 1996, 37, 3683–3692.
Tanaka, H.; Tabata, H.; Ota, K.; Kawai, T. Phys. Rev. B. 1996, 53, 14112–14116.
Tanaka, T.; Montanari, G. C.; Mulhaupt, R. IEEE T. Dielect. El. In. 2004, 11, 763–784.
Tang, H.; Ma, Z.; Zhong, J.; Yang, J.; Zhao, R.; Liu, X. Colloid Surface A. 2011, 384,
311–317.
Tang, H.; Zhong, J.; Yang, J.; Ma, Z.; Liu, X. J. Electron. Mater. 2010, 40, 141–148.
Tang, H.; Lin, Y.; Sodano, H. A. In; Goulbourne, N. C.; Ounaies, Z., Eds.; 2012; pp.
834206–834206–8.
Tchoul, M. N.; Fillery, S. P.; Koerner, H.; Drummy, L. F.; Oyerokun, F. T.; Mirau, P. A.;
Durstock, M. F.; Vaia, R. A. Chemistry of Materials 2010, 22, 1749–1759.
Tinga, W. R.; Voss, W. A. G.; Blossey, D. F. 1973, 44, 3897-3902.
Tomer, V.; Polizos, G.; Manias, E.; Randall, C. A. J. Appl. Phys. 2010, 108, 074116.
Tong, Y.; Li, Y.; Xie, F.; Ding, M. Polym. Int. 2000, 49, 1543–1547.
Traina, C. A.; Schwartz, J. Langmuir 2007, 23, 9158–9161.
Tsurumi, T.; Sekine, T.; Kakemoto, H.; Hoshina, T.; Nam, S.-M.; Yasuno, H.; Wada, S.
J. Am. Ceram. Soc. 2006, 89, 1337–1341.
Tuncer, E.; Rondinone, A. J.; Woodward, J.; Sauers, I.; James, D. R.; Ellis, A. R. Appl.
Phys. A. 2008, 94, 843–852.
Tuncer, E.; Sauers, I.; James, D. R.; Ellis, A. R.; Paranthaman, M. P.; Goyal, A.; More,
K. L. Enhancement of Dielectric Strength in Nanocomposites. Nanotechnology 2007, 18,
325704.
Urban, J. J.; Yun, W. S.; Gu, Q.; Park, H. J. Am. Chem. Soc. 2002, 124, 1186–1187.

118
Viswanathan, G.; Chakrapani, N.; Yang, H.; Wei, B.; Chung, H.; Cho, K.; Ryu, C. Y.;
Ajayan, P. M. SJ. Am. Chem. Soc. 2003, 125, 9258–9259.
Wang, F. J.; Li, W.; Xue, M. S.; Yao, J. P.; Lu, J. S. Compos. Part B.-Eng. 2011, 42, 87–
91.
Wang, X.; Zhuang, J.; Peng, Q.; Li, Y. Nature 2005, 437, 121–124.
Wang, Y.; Zhou, X.; Chen, Q.; Chu, B.; Zhang, Q. IEEE T. Dielect. El. In. 2010, 17,
1036–1042.
Wei, X.; Xu, G.; Ren, Z.; Wang, Y.; Shen, G.; Han, G. J. Cryst. Growth 2008, 310,
4132–4137.
Xie, L.; Huang, X.; Wu, C.; Jiang, P. J. Mater. Chem. 2011, 21, 5897.
Xie, S.-H.; Zhu, B.-K.; Wei, X.-Z.; Xu, Z.-K.; Xu, Y.-Y. Compos. Part A.-Appl. S. 2005,
36, 1152–1157.
Xu, C.; Xu, K.; Gu, H.; Zheng, R.; Liu, H.; Zhang, X.; Guo, Z.; Xu, B. J. Am. Chem. Soc.
2004, 126, 9938–9939.
Yang, K.; Huang, X.; Xie, L.; Wu, C.; Jiang, P.; Tanaka, T. Macromol. Rapid Comm.
2012, 33, 1921–1926.
Yang Cao; Irwin, P. C.; Younsi, K. IEEE T. Dielect. El. In. 2004, 11, 797–807.
Yen, F.-S.; Chang, C. T.; Chang, Y.-H. J. Am. Ceram. Soc. 1990, 73, 3422–3427.
Yin, W.; Kniajanski, S.; Amm, B. In; IEEE, 2010; pp. 1–5.
Yoonessi, M.; Toghiani, H.; Kingery, W. L.; Pittman, C. U. Macromolecules 2004, 37,
2511–2518.
Young, J. T.; Tsai, W. H.; Boerio, F. J. Macromolecules 1992, 25, 887–894.
Feng, Y; Yin, J.; Chen, M.; Liu, X.; Li G. In; IEEE, 2011; pp. 226–229.
Zembilgotov, A. G.; Pertsev, N. A.; Waser, R. J. Appl. Phys. 2005, 97, 114315.
Zhang, L.; Zhong, W.-L.; Wang, C.-L.; Zhang, P.-L.; Wang, Y.-G. J. Phys. D: Appl.
Phys. 1999, 32, 546–551.
Zhao, Z.; Buscaglia, V.; Viviani, M.; Buscaglia, M.; Mitoseriu, L.; Testino, A.; Nygren,
M.; Johnsson, M.; Nanni, P. Phys. Rev. B. 2004, 70, 024107.

119
Zhou, T.; Zha, J.-W.; Cui, R.-Y.; Fan, B.-H.; Yuan, J.-K.; Dang, Z.-M. ACS. Appl. Mater.
Interfaces 2011, 3, 2184–2188.

Abstract (if available)

Abstract Perovskite-based oxides like BaTiO₃, SrTiO₃, and BaₓSr₁₋ₓTiO₃ are extremely useful materials which possess fundamental properties such as high dielectric constants, high energy densities, and low loss tangents. In addition, perovskites display size, composition, and synthesis dependent dielectric properties. As such, they have found themselves at the forefront of modern capacitive and charge storage applications. While promising, such materials are typically prepared at high temperatures (>500 °C), are extremely brittle, and cannot be easily incorporated into flexible devices. In order to overcome the limitations associated with the processing of high permittivity ceramics, researchers have sought to use a nanocomposite approach, whereby small, well-defined perovskite nanocrystals are integrated into a polymer matrix. In this approach, the polymer provides processability and high breakdown strength (Ebd), while the perovskite filler delivers improved dielectric performance. With this in mind, we developed a low temperature route to preparing gram-scale quantities of small (<15 nm), well-defined BaTiO₃, SrTiO₃, and BaₓSr₁₋ₓTiO₃ nanocrystals and blended them into novel polymeric systems. Through precise control of the nanocrystal composition and surface chemistry, it was possible to better understand what factors govern the dielectric performance of the nanocrystals. Compositionally, the dielectric constant of the nanocrystal can be increased by more than an order of magnitude, while controlled surface modification enhances dispersion and improves dielectric temperature and frequency stability. This information was then used to prepare novel composite systems with enhanced performance.

Linked assets

University of Southern California Dissertations and Theses

Conceptually similar

PDF

Low temperature synthesis of functional metal oxide nanocrystals using a vapor diffusion sol-gel method

PDF

Nanocrystal surface engineering as a route towards improved photovoltaics

PDF

Expanding the library of surface ligands for semiconductor nanocrystal synthesis and photovoltaic applications

PDF

Molecular and polymeric donors for bulk heterojunction organic solar cells

PDF

Synthesis and characterization of metal chalcogenide semiconductor nanocrystals using dialkyl dichalcogenide precursors

PDF

Excited state process in perovskites

PDF

The implementation of data-driven techniques for the synthesis and optimization of colloidal inorganic nanocrystals

PDF

Quantifying the surface chemistry of semiconductor nanocrystals spanning covalent to ionic materials

PDF

Solution processing of chalcogenide functional materials using thiol–amine “alkahest” solvent systems

PDF

Surface modification of nanomaterials and developments of nanobiosensors

PDF

Molecular design for organic photovoltaics: tuning state energies and charge transfer processes in heteroaromatic chromophores

PDF

Synthesis, structural and photophysical characterization of phosphorescent three-coordinate Cu(I)-N-heterocyclic carbene complexes

PDF

Improving the efficiency and sustainability of direct arylation polymerization for conjugated polymer synthesis

PDF

Synthesis of stereoregular non-conjugated pendant electroactive polymers

PDF

Structure-function studies of random and semi-random poly(3-hexyl thiophene) copolymers

PDF

The chemistry of polymorphism in semiconductor nanocrystals

PDF

Non-conjugated pendant electroactive polymers as potential materials for opto-electronic applications

PDF

Solution‐phase synthesis and deposition of earth‐abundant metal chalcogenide semiconductors

PDF

Synthesis and characterization of 3-hexylesterthiophene based random and semi-random polymers and their use in ternary blend solar cells

PDF

Direct arylation polymerization (DARP) for the synthesis of conjugated polymers

Asset Metadata

Creator Beier, Christopher W. (author)

Core Title Synthesis and surface modification of perovskite-based nanocrystals for use in high energy density nanocomposites

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 05/21/2013

Defense Date 04/24/2013

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag barium titanate,capacitor,composite,nano,OAI-PMH Harvest,perovskite

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Brutchey, Richard L. (committee chair), Hodge, Andrea M. (committee member), Thompson, Mark E. (committee member)

Creator Email chriswbeier@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-257115

Unique identifier UC11295445

Identifier etd-BeierChris-1697.pdf (filename),usctheses-c3-257115 (legacy record id)

Legacy Identifier etd-BeierChris-1697.pdf

Dmrecord 257115

Document Type Dissertation

Format application/pdf (imt)

Rights Beier, Christopher W.

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus MC 2810, 3434 South Grand Avenue, 2nd Floor, Los Angeles, California 90089-2810, USA